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= Pt /Vd . The discharge volume Vd is given by R 2dL d where the active discharge radius is R d and the height is L d . The power absorbed by the reactor, Pt , is defined as the difference between the incident microwave power, Pinc , and the reflected power, Pref , i.e., Pt = Pinc ⫺ Pref . The discharge is usually separated from the discharge chamber walls and thus Vd < Vc . The discharge volume Vd is often adjusted (see Fig. 2) by varying the power absorbed so that deposition uniformity is
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achieved over the substrate surface. Other internal variables include neutral gas temperature, electron density and temperature, radical species concentrations, and gas residence time. 4.
Output Variables: Reactor Performance and Deposited Film Characteristics
A number of reactor process/performance figures of merit have been adopted to enable the evaluation of deposition process performance. These figures of merit, Y1 , are defined as follows: 1. 2. 3.
Linear growth rate (m/hr) is defined as the thickness increase of the diamond film (m) divided by deposition time (hr). Total growth rate (mg/hr) is defined as the weight increase of the substrate (mg) divided by deposition time (hr). Carbon conversion efficiency (CCE) is defined as the percentage of carbon atoms in the input gases that are converted into diamond. The equation for carbon conversion efficiency is CCE = Nfilm /Ninput where Nfilm is the number of carbon atoms in the deposited film and Ninput is the number of carbon atoms that entered the discharge chamber in the input gas flow.
The deposited film characteristics, Y2 , can include a wide range of properties including film uniformity, texture, morphology, thermal conductivity, electrical properties, and optical properties as well as others. The film characteristics that will be most closely tied to the microwave CVD machines in this chapter include film uniformity, film texture, and film morphology. The other film characteristics are examined in more depth in other chapters of this book. 5.
Quantifying the Input, Internal, and Output Variable Relationships
The sections that follow in this chapter quantify the relationships between the input variables and internal variables, X = f (U ), the input variables and output variables, Y = g(U ), and the internal variables and output variables, Y = h(X). In particular, Sec. IV establishes the range of reactor geometry variables, U2 , and some of the reactor performance relationships, Y1 = g(U2). Section 5 presents the output variables (both reactor performance and deposited film characteristics) as functions of the adjustable input parameters and the internal variables, i.e., Y = g(U1) and Y = h(X). Section VI describes the internal variables and their relationship to input variables based on experimental measurements X = f(U ), and lastly Sec. VII examines in more detail the internal variables based on modeling studies, X = f(U ).
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IV.
MICROWAVE PLASMA SOURCES FOR DIAMOND THIN-FILM DEPOSITION
A.
Basis for Comparison
Many different microwave plasma-assisted diamond deposition machines have been designed including tubular reactors, bell jar reactors, plasma jet reactors, ellipsoid reactors, plasma disk reactors, surface wave sustained reactors, and magneto microwave reactors. This section describes some representative reactors for the various types and gives a limited number of specifications for comparison of the reactors. The set of parameters selected for comparison includes microwave coupling structure, microwave frequency, discharge chamber size, substrate size, input microwave power range, and operating pressure range. The diamond deposition outputs (e.g., deposition rate) from each of the different reactor systems are only briefly described in this section. A detailed comparison of the diamond deposited by each system is not attempted because the different systems described here were designed on the basis of different objectives. Some of these design objectives have included designs focused on creating small inexpensive research systems, high-volume production systems, or low-substrate-temperature deposition systems. More detailed discussions of the actual diamond deposition outputs (e.g., deposition rate, film texture) will be undertaken in later sections of this chapter where the outputs will be studied as a function of input variables such as pressure, input chemistry, and flow rate.
B.
Selected Reactor Structures
1.
Tubular Reactor (NIRIM Reactor)
A schematic diagram of a typical tubular microwave reactor is shown in Fig. 5. The tubular microwave reactor was developed in the early 1980s by Kamo et al. [6] and has become the most common diamond film deposition reactor technology used for basic deposition research studies. As shown, the input gas, which is a mixture of hydrocarbon and hydrogen gas, is dissociated by the 2.45-GHz microwave energy coupled into the quartz tube inserted through a waveguide. A plunger is attached to the end of the waveguide to minimize the reflected power. The substrate is placed on a substrate holder located at the intersection of the tube and waveguide. Typical microwave power levels are 100 W to 1.5 kW and the pressure range is less than 80–100 torr. Typical quartz tube diameters are 4.5 cm or less and the substrate sizes are limited to 2–8 cm2. Tubular reactors are simple reactors to design and low in cost to construct, and as a result they have been used in many research investigations.
Microwave Plasma-Assisted Diamond Film Deposition
Figure 5
2.
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Tubular microwave plasma-assisted CVD reactor. (From Ref. 6.)
Microwave Cavity Plasma Reactor
The microwave cavity plasma reactor (MCPR) [7,8] developed at Michigan State University and Wavemat/Norton in 1986–1995 is shown in Fig. 6. Two design variations are shown including an early version (1987) in Fig. 6a and a later version (1989) in Fig. 6b. The microwave discharge is produced inside a quartz dome which is located at one end of a microwave cavity. The cavity is formed by a cylindrical wall and a movable top short defining the top end of the cavity. The microwave energy is transmitted from a 2.45-GHz microwave generator through a rectangular waveguide to a transition into a coaxial waveguide that ends as a length adjustable excitation probe in the cylindrical cavity. A hemispherical plasma discharge is created and is placed in direct contact of the substrate by adjusting the length of the probe and the location of the sliding short, i.e., L p and L s . The diameter of the microwave cavity is 17.8 cm and the diameter of the discharge region is 9–12 cm. Various substrate holders and substrate sizes can be used ranging from 5 to 10 cm in diameter. The sliding short height L s is most commonly adjusted to a length of approximately 21 cm when a 2.45-GHz microwave excitation is used. This cavity height produces a resonant
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Figure 6 Microwave cavity plasma reactor: (a) early version and (b) later version. (From Refs. 7, 8.)
microwave mode that has been identified as the TM013 mode. Typical power levels are 500 to 5000 W and pressures from 5 to 140 torr. Care is taken in the design of the gas inlet structure and substrate holder structure so that the input gas is well mixed with the discharge and to get uniform gas injection and radially uniform discharge chamber exhaust gas pumping. The substrate holder can be designed with various heights. In addition, various substrate holders have been designed that are thermally floating, cooled, and heated. Figure 6a shows an early version of the reactor, which had a discharge diameter of 9 cm and a smaller substrate diameter. Figure 6b shows an improved reactor version designed for use with up to 10-cm-diameter substrates. Also, this more recent unit operates at higher pressures up to 180 torr and powers up to 6 kW. A larger 915-MHz version of the MCPR with a microwave applicator/cavity diameter of 45 cm was also designed, constructed, and operated (1994 at Norton Diamond Film). This larger design has a discharge diameter of 33 cm and substrate diameter
Microwave Plasma-Assisted Diamond Film Deposition
Figure 6
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Continued
of up to 20 cm. It operates using 8–18 kW of microwave power generated by a 915-MHz power supply. 3.
ASTeX Bell Jar and High-Pressure Microwave Sources
The bell jar microwave plasma-assisted CVD reactor initially developed in 1987 [9,10] is shown in Fig. 7. The basic features of this reactor include a
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ASTeX bell jar microwave plasma-assisted CVD reactor. (From Ref. 9.)
silica bell jar that confines the plasma discharge. The initial designs used a 10-cm inner diameter bell jar and 7.5-cm diameter substrate holder. The microwave power generated by a magnetron supply is transmitted to the reactor using a rectangular waveguide. An antenna couples the energy from the waveguide into the reactor, which is cylindrical in shape. The position of the substrate could be moved and the substrate either was thermal floating or could be heated. The pressure range utilized was 40–70 torr and the power range for this unit was typically 1 kW. By 1992, ASTeX had improved the bell jar design and developed the ASTeX High Pressure Microwave Source (HPMS) as shown in Fig. 8 [10]. The former bell jar that confined the discharge and permitted the transmission of microwave energy was replaced by a silica microwave window. A plasma discharge was formed in the region between the silica window and the substrate holder. The position of the substrate could be adjusted to optimize the discharge-substrate interaction and the subsequent film uniformity
Microwave Plasma-Assisted Diamond Film Deposition
Figure 8
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ASTeX high-pressure microwave source. (From Ref. 10.)
on the substrate. The substrate could be either heated or cooled as needed. The pressure range in which this system operated was from 10 torr to above 120 torr. The power levels used ranged from 1 to 5 kW with later systems able to operate at powers up to 8 kW. The lack of magnetron power supplies above 8 kW at 2.45 GHz prevented the extension of this 2.45-GHz design to even larger areas at the higher pressures. ASTeX designed a 915-MHz system that permitted the use of the higher power 30–60 kW magnetron power supplies [10]. The 915-MHz system had substrate sizes as large as 30 cm. This unit was reported as depositing diamond at a rate of 10 m/hr over a 20-cm-diameter substrate. The mass deposition of diamond approached 1 g/hr. 4.
UC Berkeley Bell Jar Reactor
A schematic diagram of the UC Berkeley bell jar microwave plasma-assisted CVD system [11–14] is shown in Fig. 9. The bell jar is 10.2 cm in diameter
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Figure 9
Bell jar plasma-assisted CVD system. (From Refs. 11–14.)
with an annular gas feed system and gas exhaust system. The substrate utilized had a diameter of 2.5 cm and was supported by a quartz sample holder at the position near the plasma ball. This reactor was used extensively to study diamond deposition from solid-phase carbon sources. 5.
France-LIMHP Bell Jar Reactor
A schematic diagram of the LIMHP reactor studied by Gicquel and coworkers [15] is shown in Fig. 10. The reactor is composed of a quartz bell jar 10 cm in diameter that is surrounded by a 25-cm-diameter Faraday cage that acts as a resonant cavity for the microwave energy. The microwave power is coupled into the Faraday cage via a waveguide and an antenna. The substrate holder, which is 5 cm in diameter, can be moved up and down to obtain an optimum interaction between the plasma discharge and the
Microwave Plasma-Assisted Diamond Film Deposition
Figure 10 Ref. 15.)
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LIMHP-France bell jar plasma-assisted diamond CVD system. (From
substrate leading to good deposition uniformity. The pressure range of this system is 20–120 torr and the input microwave power range is 600– 6000 W. 6.
Ellipsoidal Reactor
A schematic diagram of the ellipsoidal reactor developed by Funer et al. [16] is shown in Fig. 11. The microwave cavity is ellipsoid in shape with the microwave power input being located at the top. The power input from the waveguide to the cavity is done using a probe antenna. The design goal in this structure was to maximize the microwave electric field strength at the location where the plasma is desired, i.e just above the substrate. Using the ellipsoidal shape maximized the fields in this region by locating the end of the antenna at one focus of the ellipsoid and the desired plasma location at the other focus. This reactor configuration was constructed at both 2.45 GHz and 915 MHz. The 2.45-GHz system used 3–6 kW of power for operation in the pressure range from 35 to 150 torr. The substrate size for this system was 5–7.5 cm. The 915-MHz system was used to deposit diamond on 5- to 15-cm diameter substrates. The power for the 915-MHz system was 20–60 kW for the pressure range 35–150 torr.
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Figure 11
7.
Ellipsoid microwave plasma reactor. (From Ref. 16.)
Surface-Wave Microwave PACVD Reactor
A diagram of a surface-wave microwave plasma-assisted reactor [17,18] is shown in Fig. 12. The microwave energy is transmitted to the reactor through a waveguide to the waveguide surfatron. The discharge is confined in a fused silica vessel. The plasma is created and stabilized by adjusting the coaxial tuning stub and the waveguide tuning stub. The reactor operates at pressures of 1–60 torr with input power on the order of 1 kW. The substrate holder is heated and made with a molybdenum susceptor. The vertical position of the substrate holder is adjustable. The reactor size is a 2.5 cm inside diameter quartz tube and the substrate size is 2 cm in diameter. The surface-wave microwave plasma-assisted reactor was also scaled to 915 MHz [19]. The substrate holder diameter used in this configuration was 8.0 cm. 8.
Microwave Plasma Jet Torch System
Figure 13 shows a microwave jet torch reactor developed by Mitsuda and coworkers [20,21] starting in the late 1980s. The unit consists of a rectangular waveguide, a microwave transition unit, and a coaxial waveguide. The coaxial waveguide consisted of a 5.72-cm-diameter outer conductor and 2.0cm-diameter center conductor. Microwave energy was generated by a 2.45GHz generator with an output of 3.8–4.2 kW. The input gas mixture used consisted of hydrocarbon, hydrogen, and argon that flowed radially into the coaxial waveguide section and then passed through a nozzle (2.2 cm in diameter) and out of the reactor. A plasma jet was generated near the nozzle
Microwave Plasma-Assisted Diamond Film Deposition
Figure 12 18.)
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Surface-wave microwave plasma-assisted CVD reactor. (From Refs. 17,
at atmospheric pressure and the resulting dissociated species arrive at the substrate. The substrate/deposition area was 2.5–5 cm2. This design, because of the high pressure, yielded linear diamond growth rates of 30 m/hr. This is one of the highest diamond deposition growth rates reported using microwave plasma excitation. 9.
Magnetomicrowave and ECR Plasma CVD Reactor
A primary motivation for the deposition of diamond at low pressures is to achieve deposition at low substrate temperatures and across large areas. By lowering the pressure, gas heating is reduced and hence the substrate heating is reduced. In addition, at lower pressures the plasma discharge size is larger for a given input power and the plasma also diffuses more freely without volume recombination. At lower pressures microwave plasma sources are often operated with the addition of magnetic fields to improve microwave power coupling to the discharge. The addition of magnetic fields of sufficient strength allows collisionless heating of the electrons by the microwave fields to occur via the electron cyclotron resonance (ECR) mechanism. The use of magnetic fields in reactors with respect to diamond deposition was summarized by Bachmann [10]. He indicated that at pressures exceeding 7 torr the electron collision rate in the plasma is too high to permit substantial
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Microwave plasma torch CVD reactor. (From Refs. 20, 21.)
ECR heating. It was also noted that the magnetic fields hardly affect the deposition results at these higher pressures. Further, it was noted that at pressures of 75 mtorr and below, where ECR starts to become important, the diamond growth is significantly affected and it becomes difficult to grow well-defined polycrystalline diamond without getting a mixture of other graphitic–amorphous carbon phases. Well-defined polycrystalline diamond films were deposited by Mantei and coworkers [22–24] using a magnetomicrowave plasma CVD reactor operating at 1–3 torr as shown in Fig. 14. This reactor was built and investigated for the low-temperature deposition of diamond films. As shown, a magnetic field was established using a 15-cm-diameter, 9 cm high NdFeB permanent magnet with a magnetic flux density of 0.42 tesla at the magnet pole face. The discharge was operated with either CW microwave energy or pulsed microwave energy. The substrates used were 2.5 cm in diameter.
Microwave Plasma-Assisted Diamond Film Deposition
Figure 14
V.
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Magneto-microwave plasma CVD reactor. (From Refs. 22–24.)
DIAMOND DEPOSITION PROCESSES/CHEMISTRIES IN MICROWAVE PLASMA-ASSISTED CVD REACTORS
This section describes the operation of microwave plasma-assisted diamond CVD reactors with respect to discharge chemistry. The focus in this section is on nongeometric input parameter variations. Whereas the previous section described reactors of various sizes, shapes, and designs, this section focuses on the general operation of microwave CVD reactors with respect to a variety of nongeometric input parameters including input gas composition, flow rate, and effect of impurities. The growth of diamond films can be divided into a two-step process beginning with substrate surface treatment or nucleation and followed by the growth process. Nucleation can be initiated through one of a number of processes including substrate surface scratching and bias-enhanced nucleation (BEN). The BEN process is usually performed in the same diamond
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deposition reactor as the diamond growth process. The BEN process consists of applying a bias voltage to the substrate to nucleate or start the diamond growth. This process is typically performed with a specific recipe that includes the bias voltage, bias time, and carbon percentage in the feed gas. The diamond deposition reactor operating conditions are then changed once the nucleation step is completed to facilitate diamond growth on the substrate. This section, as well as the rest of this chapter, focuses on the diamond growth process on a prenucleated substrate surface. Specifically, reactor design and operating conditions for bias-enhanced nucleation are not considered further in this chapter. The diamond growth process and results are discussed in this section with an emphasis on deposition done in microwave diamond CVD systems. In particular, the influence of gas chemistry, gas impurities, deposition pressure, and substrate temperature on the growth rate, quality, and texture of diamond deposited in microwave CVD systems is covered. Using the control formulation described in Sec. III.B, this section quantifies when possible the relationships between the controllable input variables, U1 , and the diamond deposition outputs, Y, for microwave diamond CVD systems.
A.
Simplified Surface Reaction/Growth Description for Microwave Plasma-Assisted CVD
The primary function of the microwave discharge is to create the radical species that participate in the growth of diamond. For example, diamond growth by using a hydrogen-methane mixture requires both atomic hydrogen and carbon-containing radicals. The microwave discharge serves to dissociate some of the molecular hydrogen into atomic hydrogen and to dissociate the methane molecules into CH3 and possible other radicals such as CH2 and CH. A general view of the deposition process [25] is one where atomic hydrogen is bonded to almost all of the surface. The primary mechanism that opens sites on the surface is abstraction of surface-terminating hydrogen by atomic hydrogen. The reaction and associated rate, Rabs⫺H , is given by CdH ⫹ H → C * d ⫹ H2 Rabs⫺H = k 1[H] where k 1 is the reaction rate constant and [H] is the atomic hydrogen concentration (flux) at the surface. The abstraction results in an open site on the diamond surface, which can be filled by either the adsorption of another atomic hydrogen or a carbon radical species (e.g., CH3). The reaction and associated rate, RadH , of hydrogen adsorption onto an open site, C d*, is
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C d* ⫹ H → CdH RadH = k 2 [H] A growth species can also be adsorbed on to the surface with the reaction and associated rate for an open surface site given by C* d ⫹ CH3 → CdCH3 RadC = k 3 [CHx ] Once the growth species is on the surface, it can proceed along one of two primary paths. The first path is thermal desorption from the surface leading to an open site on the surface again. The reaction and associated rate for a given adsorbed site are CdCH3 → C d* ⫹ CH3 R des = k 4 The other pathway is the incorporation of the adsorbed carbon species into the diamond structure. The general event that begins this mechanism in growth models is an abstraction of a hydrogen from the adsorbed carbon species. CdCH3 ⫹ H → CdCH2 ⫹ H2 Rabs⫺CHx = k 5 [H] When the species adsorbed on the surface is CH2 , the carbon is believed to change to having two carbon-carbon bonds. Subsequent hydrogen abstraction and carbon incorporations complete the incorporation of the original carbon atom into the diamond crystal. The conditions that determine the rate of the five mechanisms listed above are surface temperature, gas-phase atomic hydrogen concentration at the surface, and gas-phase carbon radical concentration at the surface. In steady-state conditions, this set of reactions can be combined to give a growth rate G given by [25,26] G = k3
ns nd
冉
k1 k1 ⫹ k2
冊
[CH3][H] k4 ⫹ [H] k5
(1)
where ns is the surface site density (2.61 ⫻ 10⫺9 mol/cm2 on (100) surfaces) and nd is the molar density of diamond (0.2939 mol/cm3). During the deposition process, the surface is understood to be covered nearly 100% with atomic hydrogen as carbon radicals are being deposited in the diamond (sp3) phase. Low hydrogen surface coverage results in sp2like terminations on the diamond surface, which lead to defects and if the
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hydrogen coverage is low enough graphitic/amorphous film deposition. An empirical model [25] for the defects, Xdef , incorporated during diamond growth is Xdef ⬀
G [H]n
where n is an empirical fit that is often selected at n = 2. For n = 2 the simplest defect concentration model is given by Xdef ⬀
[CH4 ] [H2 ]
Hence, the higher quality (lowest defect density) diamond is deposited when the carbon species concentration in the input gas flow (e.g., methane concentration/hydrogen ratio) is low. Other work [4,27] has established the condition for achieving diamond growth. One necessary condition is to produce a flux of atomic hydrogen on to the surface that exceeds the carbon incorporation rate into the film by a factor of approximately 7000–11,000. Hence a key function of the plasma discharge in microwave CVD reactors is to provide a flux of atomic hydrogen that exceeds the carbon incorporation rate by a factor in the range of 104. To summarize, the many processes occurring on the diamond surface include adsorption of atomic hydrogen and carbon-containing species, hydrogen abstraction from the surface, and desorption of carbon radicals from the surface. The important variables for the diamond deposition process, including the microwave plasma-assisted CVD systems being considered in this chapter, are (1) substrate temperature, (2) atomic hydrogen flux, (3) carbon radical flux and composition, and (4) other ‘‘impurity’’ species fluxes. The rest of this chapter focuses on the discharge chemistry and diamond deposition results for microwave plasma-assisted CVD systems. Microwave plasma-assisted diamond CVD reactor performance as described earlier in Sec. III.B can be looked at either as the input variable settings determining the film properties (the approach used in this section, Sec. V, of this chapter) or as the input variable settings determining the reactor internal variables (e.g., [H], [CH3 ]) and then the internal variables determining the film properties (the approach used in Secs. VI and VII of this chapter). A somewhat reduced process variable model as shown in Fig. 15 is used as the basis for the discussion in the rest of this section. The input variables are input reactant gas composition, input gas flow rate, substrate temperature, input microwave power, and pressure. The output variables to be assessed include growth rate, diamond film quality, and film texture. The approach in this section is to give a general discussion of the trends for diamond growth in microwave CVD systems for a pressure range
Microwave Plasma-Assisted Diamond Film Deposition
Figure 15 process.
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Reduced variable model of the microwave plasma-assisted CVD
from 1 to 200 torr for a variety of input gas chemistries. The key reactor input parameters and process quantities to achieve diamond deposition uniformly across a substrate surface are as follows: 1. 2. 3. 4.
5. 6.
Establish the desired input gas flow composition and rate. Establish and maintain a uniform and selected temperature across the deposition surface of the substrate. Establish the exit gas pumping rate to get the desired pressure. At higher pressures construct the input gas flow, discharge chamber, and exit gas flow to optimize the flow dynamics for uniform diamond deposition. Construct the reactor to prevent or minimize contamination of the plasma discharge from leaks and reactor materials via wall erosion. Input the microwave power to establish the plasma discharge. Optimize the input power magnitude, discharge chamber size and cooling, microwave power absorption profile, and substrate position to get uniform deposition across the substrate area.
The first generality in understanding the relationship of the deposition chemistry and output variables is the region of diamond growth with respect to input gas composition. The most commonly used input gas mixtures consist of hydrogen-, oxygen-, and carbon-containing molecules. Because the temperatures and activation of the gases in the plasma discharge are extensive, a primary determining factor for the film composition is the C, H, and O atomic molar composition of the input gas, regardless of the exact input
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species molecular composition. A useful summary of the input gas composition appropriate to grow diamond was developed by Bachman et al. in 1990 [1] and further refined in 1994 [28]. The plot in Fig. 16 shows such a ternary C-H-O diagram. The carbon deposition region exists very near and just above the H-CO line connecting 100% H with C-50% and O-50%, i.e., the crosshatched region shown in Fig. 16. The diamond growth region is very narrow to nonexistent on the right end of the H-CO line where [H] approaches zero and becomes wider on the left side of the H-CO line when [H] is larger. Once a specific input gas composition is selected in Fig. 16, the diamond deposition is still a function of a number of other variables. Several other variables are important including pressure, flow rate, and substrate temperature, as described in more detail in the following subsections.
B.
Methane-Hydrogen Deposition Discharge
It is useful first to summarize the general trends for diamond deposition using the simplest discharge chemistry consisting of methane and hydrogen. To do this, some background needs to be developed first. The diamond growth proceeds in a methane-hydrogen discharge as described earlier in this section primarily to supply the [H] flux to the substrate surface and carbon radical flux to the surface. The most important carbon flux in the microwave plasma considered here is the CH3 methyl radical. To deposit diamond, the [H] flux to the surface is quantified as being 7000 to 10,000 times larger than the carbon flux that is incorporated in the growing film [4,27]. When the [H] is larger than this factor, diamond is deposited on a diamond surface. Further, the larger the [H]/[CH3 ] ratio is at the growth surface, the less sp2-bonded carbon and defects are incorporated into the film and the higher is the diamond quality. The other general observation is that the growth rate is proportional to the [CH3 ] flux to the surface as indicated in Eq. (1). The general trends for the growth rate versus temperature, pressure, and methane percentage in microwave plasma-assisted diamond CVD systems are shown in Figs. 17, 18, and 19 respectively. The general trends are 1. 2.
3.
An increase in the growth rate as the pressure is increased [29] as shown in Fig. 17. An increase in the growth rate as the temperature increases up to a maximum, then a drop in the growth rate at higher temperatures [5,30] as shown in Fig. 18. An increase in the growth rate that is linear versus methane percentage at low methane percentages, that is sublinear versus methane concentration for moderate methane percentages, and that
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Figure 16 C/H/O diagram of diamond CVD developed by Bachmann et al. (From Refs. 1, 28.)
saturates versus methane concentration at higher methane concentrations. These trends can be seen in the experimental data shown in Fig. 19. After the saturation value is reached, the growth rate may drop slightly as the methane concentration is increased further. Typically, in this high methane concentration region, nondiamond carbon or diamond with extensive secondary nucleation is predominantly deposited [5,26]. 1.
Growth Rate versus Pressure and Methane Percentage
One objective of this chapter is to quantify some of the experimental data obtained for microwave CVD systems across a wide parameter space. An empirical model for the growth rate G is G ⬃ R␥ where R is the methane fraction. In work by Harris and Weiner [31] the value of ␥ was pressure
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Figure 17 The total weight gain and absorbed microwave power density versus pressure. Note that the data in the 30–80 torr range include 7.5-cm and 10-cmdiameter substrates. The data above 80 torr are for 5-cm-diameter substrates. (From Ref. 8.)
dependent, varying from ␥ = 1 at 20 torr to ␥ = 0.5 at 200 torr. This growth rate model was developed in a hot-filament deposition system. A similar model can be developed for the microwave deposition system. To develop such a model for the microwave system we will consider data across a wide range of pressure and methane concentrations. Some data showing deposition rate versus methane percentage and pressure are shown in Fig. 20 for pressures of 18, 25, 53, and 135 torr [5,26,30,32]. The temperature used for these growth rates is at or near the measured or generally expected temperature range where the growth rate peaks. The methane concentration (#/unit volume) in the plasma discharge is a function of the discharge pressure, the discharge temperature, and the methane percentage in the input gas flow. In a microwave plasma source the discharge gas temperature increases as the pressure increases (see Sec. VI.B), assuming the discharge volume stays approximately constant. The neutral density in the discharge gas is proportional to (pressure)/(gas temperature). An approximate fit to the neutral gas
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Figure 18 Diamond growth rate as a function of deposition temperature and methane concentration. (From Ref. 30.)
density versus pressure is given by density ⬃ (pressure)0.75 (see data in Sec. VI.B). Hence the density does not increase linearly with the pressure because of the gas heating that occurs as the pressure is increased. This is an empirical fit that models the discharge neutral density versus pressure. The full model for the growth rate is GTsmax = Cp0.75 R where C is a proportionality constant, p is the pressure, and R is the methane percentage. Figure 20 shows a plot of this model and experimental data for growth rates in microwave CVD system for deposition pressures from 18 to 135 torr. This model is for the region of methane concentrations where the growth rate varies linearly with R. The growth rate increases linearly versus the methane concentration for low methane concentrations, but increases beyond this linear region are characterized by the growth rate changing sublinearly versus methane increases as shown in Fig. 19. This can be understood by looking deeper at the carbon chemistry in the plasma discharge. The linear region at low methane concentrations is characterized by the growth species in the plasma
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Figure 19 Linear growth rate of diamond versus methane concentration. (From Refs. 5 (a) and 26 (b).)
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Figure 20 Comparison of growth rate model with experimental data from Refs. 5, 26, 30, and 32. The growth rate model equation is (growth rate = Cp0.75R). This model is valid for low methane concentrations and for a range of substrate temperatures near the substrate temperature that gives a maximum deposition rate.
(CH3) increasing linearly with the methane concentration. Once the methane concentration becomes large enough, the CH3 concentration no longer increases linearly with the methane concentration as shown in Fig. 21 [26]. The gas-phase chemistry present in the discharge converts much of the CH4 to C2H2. This boundary where the CH3 percentage changes from a linear relationship versus methane concentration to a sublinear relationship varies as a function of pressure. A plot of the region of linearity versus pressure is shown in Fig. 22 based on experimental data [5,26,30,32]. The region below
Figure 21 CH3 and atomic hydrogen concentration versus input methane percentage as modeled by Farhat et al. [26] for a pressure of 7000 Pa and a plasma gas temperature of 2850 K.
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Figure 22 Regions of growth rate versus methane percentage as a function of methane percentage and pressure based on the experimental data of Refs. 5, 26, 30, and 32.
the dashed line exhibits a linear growth versus methane concentration. This plot uses data for substrate temperatures in the range where the maximum deposition rate occurs versus substrate temperature. 2.
Quality versus Pressure and Methane Percentage
The quality of diamond films can be defined in a number of ways including the number of incorporated defects and the Raman signal. Quality can also be defined in terms of the intended application for the diamond. Some measures of quality include the number of defects incorporated in the diamond crystals, the number or density/rate of secondary nucleation sites, the size of individual crystals in the polycrystalline film, and the nature of the grain boundaries. The general trend is that the larger the hydrogen flux is to the surface and the slower the growth rate (carbon incorporation into the growing film), the higher is the quality. In terms of the Bachmann diagram, the closer the plasma composition moves toward the C-O line, i.e., toward the no-growth region, the higher the quality of the deposited film. Because the definition of diamond film quality often depends on the intended application of the film, a simple definition of quality will be adopted here. This definition of quality has two regimes: the higher quality deposition regime, which has well-defined crystals, and the lower quality regime, which does not have well-defined crystals; i.e., it is extensively renucleated and often referred to as cauliflower-like films. Alternatively, a quality transition can be defined as the transition from a low secondary nucleation value, where well-defined crystals are formed, to a high secondary nucleation value, where cauliflower-like growth occurs. This region of
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high-quality growth in a methane percentage and pressure parameter space is shown in Fig. 22 based on experimental data [5,26,30,32]. It is the region below the solid line in the figure. The cauliflower diamond growth region is above the solid line in Fig. 22. 3.
Growth Rate versus Substrate Temperature
The deposition rate of diamond from methane-hydrogen gas mixtures generally shows a maximum rate versus substrate temperature as seen earlier in Fig. 18. The temperature range of this peak is in the range 850–1150⬚C with the location of the maximum deposition rate being dependent on the pressure. Experimental results [5,26,30,32] have shown that the maximum growth rate temperature, Tmax , increases gradually with pressure as shown in Fig. 23. Some specific data for the growth rate versus substrate temperature that yields the maximum growth rate include work by Gicquel et al. [30] at 18 torr that gave peak values for substrate temperatures of 850– 950⬚C, Kondoh et al. [33] at 30 torr that gave peak values for substrate temperatures of 900–1000⬚C and Kuo [5,29] at 135 torr that gave peak values for substrate temperatures of 1050–1150⬚C. In even higher pressure, nonmicrowave CVD processes, the substrate temperature of maximum deposition rate at atmospheric pressure is 1200⬚C in work by Snail and Marks [34]. At temperatures less than Tmax , the growth rate dependence on substrate temperature is generally modeled with an Arrhenius expression using an associated activation energy Ea . The general growth rate expression is G ⬀ exp
冉 冊 ⫺Ea RT
Figure 23 Substrate temperature for maximum growth rate, Tmax , i.e., the shaded region, versus pressure from Refs. 5, 26, 30, and 32.
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where G is the growth rate, R is the gas constant, and T is the temperature in kelvin. In the temperature range 440–510⬚C, Ulczynski et al. [35] reported an activation energy of 14.6 kcal/mol in the pressure range 5–15 torr. Kondoh et al. [33] using a hot-filament system reported an activation energy of 23 kcal/mol. The data of Gicquel et al. [30] in the temperature range 700–850⬚C showed an activation energy in the range 11–16 kcal/mol. In fact, many results from growth rate measurements of polycrystalline diamond deposition show activation energies of 15–30 kcal/mol [4]. 4.
Growth Rate versus Flow Rate
A key parameter determining the importance of the flow rate is the residence time of the feed gas in the deposition chamber. The residence time, t, is given by t = V/F where V is the discharge chamber volume and F is the flow rate scaled to the pressure in the chamber. Example residence times in a typical discharge chamber range from 1 sec to a few minutes. For example, a 200 sccm (STP) flow is 7600 cc/min at 20 torr and 1000 cc/min at 150 torr. So if the discharge chamber volume is 1 L (1000 cm3) the gas residence time is about 8 sec at 20 torr and 60 sec at 150 torr. While the feed gas is within the plasma discharge region, it is undergoing chemical reactions and some of the carbon is being deposited as diamond on the substrate. A measure of how much of the carbon in the input feed gas is deposited on the diamond substrate is the carbon conversion efficiency. The carbon conversion efficiency increases as the flow rate is reduced as shown for some experimental data in Fig. 24 [29,36,37]. In terms of the effect of this carbon consumption (or sink) on the gasphase chemistry, lower flow rates lead to a gas-phase composition of the discharge that has a lower carbon content than present at higher flow rates. For example, if the input flow has a carbon input of 0.01 mole fraction and the carbon conversion efficiency is 25%, then the exit gas has a carbon concentration of 0.0075 mole fraction. Because of the long gas residence time in typical discharge chambers for diamond growth, the exit gas composition can be considered a close approximation of the carbon mole fraction in the chamber. Hence, to first order a lower flow rate is equivalent to a lower carbon concentration. This is particularly true when the carbon conversion efficiency is high. However, the possible anticipated higher quality diamond expected from lower carbon concentrations (i.e., lower flow rates/ higher carbon conversions) is often not observed because another effect of
> Figure 24
Carbon conversion efficiency versus flow rate. (From Refs. 36, 37.)
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lower flow rates and the associated longer gas residence times is that any leaks or chamber wall erosion will lead to larger impurity concentrations in the gas discharge. Such impurities can alter and sometimes degrade the deposited film process and properties. An example of the influence of flow rate is the work by Ralchenko et al. [37]. A microwave system using methane and hydrogen at 100 torr, 2.45 GHz and a substrate temperature of 720⬚C was utilized to grow diamond at flow rates ranging from 60 to 1000 sccm. They found the highest quality diamond films intended for millimeter wave windows were deposited at the higher flow rates. This result is suggestive of some impurity in the deposition process being reduced by the higher flow rate. 5.
Growth Rate versus Deposition Time
The growth rate as a function of the time is not constant for the entire duration of the film growth process. Initially the growth rate is slow as the film first nucleates and grows to cover the entire surface of the substrate. Next the growth rate increases as the film texture is formed and the crystals oriented with their fastest growth direction perpendicular to the substrate surface outgrow the other less favorable crystal orientations. Finally, once the film has become well textured, the growth rate becomes constant versus time. This phenomenon is indicated in Fig. 25a for growth at a pressure of 135 torr for 120 h [5]. The same effect was demonstrated in the work of Ralchenko et al. [37] at a pressure of 100 torr as shown in Fig. 25b. 6.
Dependence of Film Texture and Growth Parameter ␣ on Discharge Chemistry
As diamond grows, the shape of the crystals depends on the growth rate of the various surfaces of the crystals. A useful measure for understanding the relationship of growth on various surfaces is to study the growth along some of the primary directions in the diamond crystal structure. One finds that the growth rate along the three directions [100], [110], and [111] changes as a function of gas chemistry and substrate temperature [38]. The growth rate along the [110] direction is generally the fastest. As a result the {110} surfaces are rarely seen in the polycrystalline films grown in CVD reactors, including microwave reactors. This occurs because any (110) surfaces will grow out quickly, leaving other surfaces with slower growth rates defining the crystal. The growth rate in either the [100] direction or the [111] direction is next fastest depending on the growth conditions. Maeda et al. [39] did a study of the growth rate along these two directions in a microwave CVD
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Figure 25 (a) Linear growth rate of as a function of deposition time [5]. (b) Growth rate versus film thickness. (From Ref. 37.)
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reactor. The study (Fig. 26) was done so that the growth rates along the [100] and [111] directions were measured as a function of methane concentration and substrate temperature. A useful parameter for understanding how the growth rates along the various directions determine the crystal shape is the growth parameter ␣ defined as
␣ = 兹3
V100 V111
where the growth rates along the [100] and [111] directions are given by V100 and V111 , respectively. When ␣ = 1 or less, the [111] direction grows quickly and the individual isolated crystals formed are cubes with (100) surfaces as shown in Fig. 27. Fig. 27 shows the diamond crystals that grow for values of ␣ ranging from ␣ = 1 to ␣ = 3 [40]. Also shown in Fig. 27 is the direction of fastest crystal growth for the various ␣ parameter values. The direction of fastest growth is indicated by the arrow by each cubooctahedral shape. For ␣ = 3 or larger, the [100] direction grows most quickly and the crystal shape formed has (111) surfaces. Conversely, for ␣ = 1 or less, the [111] direction grows most quickly and the crystal shape is formed by the more slowly growing (100) faces. In the case of ␣ = 1 or less, the longest direction in the crystal is the [111] direction. For ␣ = 3 or larger, the longest direction in the crystal is the [100] direction. At an intermediate value of ␣ = 1.5 the longest direction in the crystal is [110]. A plot of the ␣ parameter versus methane concentration and substrate temperature from the work of Maeda et al. [39] is shown in Fig. 28. The picture becomes more complicated when a polycrystalline film is grown because many crystals intersect and overgrowth occurs. The growth rates along the [100] and [111] directions determine the texture of the film and the surface morphology of the film. The texture of the film is the crystal direction that is longest in the individual crystals of the polycrystalline film. The texture in polycrystalline diamond films develops because of the crystals which have an orientation that has a direction of fastest growth perpendicular to the substrate surface. These crystals overgrow the crystals with other orientations. This is shown in Fig. 29, where a random orientation of small crystals is shown at the bottom of the film. As the film grows, the crystals that grow with their direction of fastest growth perpendicular to the surface become taller than their competitor crystals. Ultimately, these crystals with their fast growth direction normal to the substrate in Fig. 29 overgrow the competition, ending the possibility of further growth of the other crystals with less preferred orientations. A reliable measure of the ␣ parameter can be obtained using a texture analysis of the polycrystalline films from x-ray diffraction measurements.
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Figure 26 Growth rate along the [100] direction (a) and along the [111] direction (b) versus substrate temperature for various methane concentrations. (From Ref. 39.)
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Figure 27 The cubo-octahedral shapes that result as the ␣ parameter is varied from 1.04 to 2.85. These are the shapes for isolated crystallites. (From Ref. 40.)
Measurements of the velocity of growth of diamond along various crystalline directions in microwave CVD systems have been made in work by Wild, Koidl, and their colleagues [41,42] as shown in Fig. 30 and by Maeda et al. [39] as shown in Figs. 26 and 28. The trend shown in these studies is that as the temperature increases the ␣ factor decreases. Also, these studies in general showed that as the methane percentage is increased, the ␣ factor increases. These figures display the general trends for ␣ only. The exact value of ␣ at a specific substrate temperature and methane percentage may change from reactor system to reactor system because different reactor geometries result in a different relationship between the input, internal, and output variables; i.e., the different reactor geometries discussed in Sec. IV have unique and different Pt , p, and Ts operating curves as discussed earlier in Sec. III.A.3. As will be noted in the following sections on the influence of various impurities, the addition of impurities also affects the growth along various crystalline directions and consequently the ␣ parameter value. The other property of polycrystalline diamond films strongly dependent on the growth rates in various directions is the surface morphology. The surface morphology can be defined as the crystal shapes, orientations, and
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Figure 28 Growth rate ratio ([100]/[111] growth rate ratio) versus substrate temperature for various methane concentrations. (From Ref. 39.)
sizes as observed on the top surface of the growing polycrystalline film. Relating the direction of maximum growth rate and ␣ with the surface morphology of polycrystalline films must be done with great care. To appreciate this relationship, consider some randomly oriented seed crystals grown with ␣ = 1. As the individual crystals grow, they will have a cubic shape similar to the shape shown for ␣ = 1 in Fig. 27. Once the cubes start growing together, the surface will initially have the appearance of many cubes all of random orientation with respect to each other and the substrate surface. As the film grows to a larger thickness, the 具111典 texture forms because for ␣ = 1 the [111] direction grows fastest (see the arrow direction for the ␣ = 1 shape in Fig. 27). So as the film grows perpendicular to the substrate surface, the [111] direction will dominate and clear cube shapes will no longer be clearly seen. Rather, the corners of the cubes formed by overlapping and overgrowing crystals will be seen on the top surface. As a general statement, what the surface of films thick enough to have developed texture looks like can be roughly estimated as that seen by looking back down along the direction of the arrow for the crystals in Fig. 27. Commonly, films with square {100} facets are desired, and these can be formed in two basic ways. The first way is to operate with ␣ near 1. This will produce the square shapes during the initial stages of growth when the film thickness is only on the order of the initial spacing between the seed
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Figure 29 Sketch of polycrystalline diamond film growth from isolated, randomly oriented crystals shown at the bottom of the figure. The dashed lines indicate the grain boundaries. The solid lines show the surface of the film at selected times. (From Ref. 40.)
crystals or nucleation sites. Specifically, randomly oriented seed crystals will grow into small cube shapes with ␣ = 1. Continued growth under low secondary nucleation conditions will form a 具111典 texture because this direction grows the longest dimension (fastest). So observing square (100) facets when the films are thin may indicate values near ␣ = 1. The square facets in this case have orientations that are not well aligned with the substrate surface, i.e., faces parallel to the growth surface. Another way to form square facets is to grow a thick film with a 具100典 texture. This is done by using ␣ value slightly less than 3. As the film grows thick, the texture forms and the [100] directions are perpendicular to the substrate surface. The value of ␣ slightly less than 3 leads to the square tops on the crystals in the film. See, for example, the square tops on the crystal shapes for ␣ = 2.0 and ␣ = 2.25 in Fig. 27. These square facets are well aligned parallel to the surface of the growing film because of the strong texture of the film. This is the growth condition that leads to smooth coplanar (100) faceted films. So care must be used in interpreting the ␣ parameter value from just the examination of the surface morphology of polycrystalline films.
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Figure 30 Growth parameter (␣) versus methane concentration for three substrate temperatures. An ␣ parameter of 1.5 gives a 具110典 texture and an ␣ parameter of 3.0 gives a 具100典 texture. (From Refs. 41, 42.)
Once the ␣ parameter has been mapped, two-step processes [43] have been developed that use an ␣ = 3 or near 3 to grow highly 具100典 textured films during the initial growth phase. This gives the film structure shown in Figs. 31a and 32b. Then for the second growth phase the growth conditions are changed so that the ␣ parameter is reduced in value. This increases the relative growth rate of the [111] direction relative to the [100] direction and the film growth fills in the valleys located between the (100) faces present in the initial growth phase. The resulting film has a top surface covered with (100) facets all reasonably parallel with the substrate surface as shown in Figs. 31b and 32b. Another factor that influences the surface morphology is the extent and influence of twinning that can occur. Wild and coworkers [42] indicated that the ␣ parameter also influences the formation and effect of twinning. They found that for ␣ > 2 the {100} facets are stable with respect to twinning.
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Figure 31 Schematic drawing of a two-step diamond deposition process. Step one grows the film so that is has a 具100典 texture, then step two grows the film so that smooth (100) faces dominate the top surface of the film. (From Ref. 43.)
Stable means that if a small twin occurs on a (100) face it will disappear as the film continues to grow. Conversely, an unstable condition is one where a small twin grows larger as the growth continues. The other stability condition noted in their work was that for ␣ < 1.5 the {111} facets are stable with respect to twins with an inclined twin plane. C.
Influence of Oxygen (Hydrogen-Methane-Oxygen Discharges)
The addition of oxygen to the input gas flow, either as molecular oxygen, O2 , or as another gas such as CO or CO2 , can be useful for diamond deposition. The use of some oxygen in the input gas (either O2 , CO, or CO2) generally yields films with larger grain sizes and less sp2 material. Also, the inclusion of oxygen has been shown to increase the deposition rate in some cases and the addition of oxygen allows deposition to occur at lower substrate temperatures.
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Figure 32 SEM pictures of a polycrystalline diamond film after step one (a) and after step two (b and c) as described in the caption of Fig. 31. (From Ref. 43.)
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A basic requirement to achieve diamond deposition is that the C/O ratio must be near one (more specifically, equal to or slightly greater than one as shown earlier in the crosshatched region of Fig. 16). For oxygen-rich discharges no diamond growth occurs and for carbon-rich discharges (except those with large atomic hydrogen concentrations) nondiamond carbon is deposited. Any excess oxygen will etch the carbon from the surface, preventing growth, and any excessive carbon in large amounts is deposited as nondiamond carbon films. In general, a trade-off exists between higher growth rates (and the associated lower quality) that occur on having slightly more carbon versus lower growth rates (and the associated higher quality) that occur at slightly higher oxygen concentrations. Some additional understanding of diamond deposition with oxygencontaining discharges can be achieved by considering the gas-phase chemistry. The highly activated, high-temperature discharge environment of diamond deposition drives the mixture of carbon and oxygen species in the reactor toward primarily a CO molecular composition. If the feed gas mixture has substantial percentages of oxygen and carbon present, most of the oxygen and carbon are converted into CO. Four statements can be made regarding the influence of oxygen on the diamond deposition process: 1. 2.
In general, with the addition of oxygen, the substrate temperature can be dropped 100–150⬚C without a loss of growth rate [10]. The use of oxygen changes the growth rate, under some conditions to a higher rate. Figure 33 shows the growth rate versus percent oxygen (O2) added in the feed gas flow [44]. The deposition conditions used in this study by Weimer et al. [44] consisted of a tubular reactor operating at a 21 torr pressure, 170 W microwave power, and a 97 sccm H2 flow rate. The top portion of the figure shows the growth rate with a 2 sccm methane flow. It can be seen that as the oxygen flow is increased from 0 to 0.06% in the feed gas, the growth rate increases by a factor of 1.6. A similar phenomenon occurs for an acetylene feed gas as shown in the lower portion of Fig. 33. Also observed is that as the oxygen percentage increases further the growth rate drops. This presumably occurs because as the oxygen content in the discharge is increased more CO is generated, leaving less active carbon species available for deposition. Further, as the feed oxygen percentage increases, more oxygen radicals are available to etch the surface. In another study [10] of the influence of oxygen on the growth rate, the gas-phase chemistries near C/O/H = 1:1:1 gave a maximum deposition rate. This can be seen in Fig. 34, where the deposition rate peaks at the C/O/H = 1:1:1 condition [10].
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Figure 33 Diamond growth rate versus oxygen percentage in the feed gas flow. The hydrogen gas flow rate was 97 sccm. (From Ref. 44.)
3.
The oxygen reduces sp2 incorporation in the film and also reduces secondary nucleation, which leads to larger crystals and altered surface morphology. This is believed to occur because of the increased etching effect of oxygen radicals on graphitic structure in the film [45]. It has also been noted by Harker and DeNatale [46] that the crystallite size increases as increasing oxygen content is
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Figure 34 Diamond growth rate versus C/H/O feed gas composition for two pressures of 70 and 140 torr. The feed gas composition of C/H/O = 1:1:1 gave the highest growth rate. (From Ref. 10.)
4.
added to the gas feed due to increased reactive etching rates during deposition that result in decreased crystal defect density. The introduction of oxygen into the feed gas influences the ␣ parameter (growth rates along the [100] and [111] directions), film texture, and film morphology. Figure 35 shows a plot of the parameter ␣ versus methane concentration at a fixed substrate temperature of 800⬚C [42]. The first of two observations to be made is that when oxygen is present in the discharge, variations in the methane concentration produce changes in the ␣ parameter in a manner similar to the no-oxygen case (i.e., ␣ increases as methane concentration increases). The second observation is that oxygen introduction with other conditions held constant produces shifts in the ␣ parameter and hence the texture and surface morphology. The change of ␣ (as indicated by film texture) with changes in the oxygen input flow was also demonstrated by Gu et al. [47]. In their study CO (0–10%) was added to a 100 sccm H2 and 1 sccm CH4 feed gas flow. The deposited films showed a large increase in the (100) XRD peak (more 具100典 texture) as the CO flow increased from 0 to 1%. Further, as the CO flow percentage in-
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Figure 35 Growth parameter ␣ versus methane concentration for two different carbon monoxide concentrations of 17 and 67%. The substrate temperature was constant at 800⬚C. (From Ref. 42.)
creased to 4% and 10% the (100) XRD peak decreased, indicating a reduction in the 具100典 texture. A full mapping of the ␣ parameter dependence on oxygen content, hydrogen content, carbon content, and substrate temperature has not been done [40].
D.
Influence of Nitrogen Gas on Diamond Synthesis
In the mid-1990s it was demonstrated experimentally that the addition of small amounts of nitrogen (N2) gas to the conventional CH4-H2 diamond film synthesis gas has an important impact on the CVD synthesis of diamond films [48–53]. For example [49], high N concentrations result in nanocrystalline and black films containing graphite or amorphous carbon, and very high N concentrations, i.e. 20% N concentration can completely inhibit diamond growth. In contrast, the addition of only 20–100 ppm of nitrogen can have a beneficial effect. At a deposition pressure of 20–60 torr, an N
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concentration of only a few tens of ppm stabilizes the formation of smooth, 具100典 textured diamond films. In particular, Locher et al. [48] using a tubular reactor demonstrated that at intermediate methane concentrations (1–2%) at 37 torr, the addition of several tens of ppm of N2 stabilized the formation of smooth 具100典 textured diamond films. When no nitrogen was added with the CH4 percentage held equal to 1.5%, the film deposited had fine grains indicative of a high secondary nucleation rate or high twinning rate. When the nitrogen input level has set to 60 ppm, well-defined (100) faceted crystals grew with 具100典 texture. The introduction of nitrogen also increases the growth rate along the [100] direction which shifts the ␣ parameter to a higher value. This was also shown in the tubular reactor experiments of Jin and Moustakas [49]. The texture of the films shifts from 具110典 to 具100典 when tens of ppm to a few hundred ppm of nitrogen is added. This is further demonstrated in a homoepitaxial study by Wild et al. [50], who showed that the addition of 10–50 ppm of nitrogen increased the [100] growth rate much more than the [111] growth rate. Hence, the addition of small amounts of N2 to the deposition process can be used to stabilize the 具100典 growth surface. At higher deposition pressures and higher input power densities, the deposition rate was observed to increase by a factor of 3–5. Muller-Sebert et al. [51], employing a tubular reactor that operated at a pressure of 167 torr, a deposition temperature 830⬚C, and a methane concentration of 3%, demonstrated that the addition of only 25 ppm of nitrogen changed the growth rate from 2.6 to 13 m/hr. Asmussen et al. and Ayres et al. [52], using a microwave cavity plasma reactor as shown in Fig. 6b, also studied the influence of nitrogen concentration from 5 to 1000 ppm at pressures of 38 and 50 torr. They noted a variation in growth rate of less than 10% over the range 5–1000 ppm. The substrate temperatures used in this study were T = 850⬚C at 38 torr and T = 910⬚C at 50 torr. Examples of how the nitrogen addition influenced the surface morphology are shown in Fig. 36. These experiments indicate that, as in the tubular reactor experiments, the addition of small amounts (200– 400 ppm) of nitrogen stabilizes the growth of high-quality 具100典 faceted films. However, the actual threshold nitrogen concentrations and the variation in these threshold concentrations versus the other independent experimental variables differ considerably from the tubular reactor performance, suggesting that reactor geometry has an important impact on the deposition process in the presence of nitrogen impurities. It is also interesting to note that even though nitrogen has a strong effect on the growth process, only a very small amount of the nitrogen is actually incorporated into the film [49,53]. Typical values for the incorpo-
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Figure 36 SEM images of nine samples deposited with 200 sccm flow rate, 38 torr pressure, and 1% methane. Nitrogen gas-phase concentration and diamond Raman signal peak width are given at the top of each micrograph. (From Ref. 52.)
ration probability are 2 ⫻ 10⫺4 for the {100} faces and 7 ⫻ 10⫺4 for the {111} faces [53]. E.
Chlorine- and Fluorine-Based Deposition Chemistries
Diamond has been deposited with halogen gases and halogen-containing gases included in the feed gas flow. Halogen-containing gas mixtures including CF4 , Cl2 , CH3F, CH3Cl, C2H2Cl2 , C2H5Cl, C3F8, and others have been used to grow diamond [10,54]. Some of the advantages of using hal-
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ogens as noted in the literature include enhancement of growth rate, improvement of diamond quality, and reduction of growth temperature. The role of chlorine or fluorine in the deposition process is believed to be due to hydrogen abstraction from the surface. A model [55] for fluorine and chlorine is that the abstraction of the hydrogen from the surface is done more quickly by atomic fluorine and atomic chlorine at lower temperatures. This opens up sites more often for the attachment of a carbon species (such as CH3). The region of diamond growth for the C/H/Cl system in a microwave CVD system [10,56] occurs for a mixture of predominantly hydrogen with carbon composition percentages of 2–3% or less and chlorine composition percentages of 40% or less to get diamond. The region of diamond growth for the C/H/F gas composition was again for a mixture of predominantly hydrogen with C percentages of 2% or less and a fluorine composition percentage of 12% or less. The advantage of using halogen to lower the deposition temperature in hot-filament systems is well documented. Such an advantage is less clear in microwave CVD systems. Bachmann et al. [10,56] reported studies of C/H/F and C/H/Cl systems in which no distinct advantage in terms of growth rate or lower deposition temperature was found as compared with the C/H/O system. F.
Inert Gas/Methane/Hydrogen Deposition Chemistries
Two basic influences of adding a noble or inert gas to the input gas flow for diamond deposition have been identified. The first is an enhancement of the atomic hydrogen and the growth species CHx concentration in the plasma discharge by the addition of Xe. This work was done by Hosomi and coworkers [57]. They observed an increase in the deposition rate of diamond on adding 1% Xe to the feed gas composition of 1 part methane and 99 parts hydrogen. They used a microwave CVD system operating at 800 W 2.45 GHz, and 40 torr. The low excitation energy level of Xe resulted in a plasma discharge with a higher electron density. This then produced more atomic hydrogen and CH3 and a 50% increase in the deposition rate with the Xe addition. The second influence of inert gases occurs when large quantities of a noble gas are added to the input gas flow. In a series of studies [58,59] at Argonne National Laboratory, large percentages of argon in the input gas flow together with low hydrogen percentages have been used to deposit nanocrystalline diamond. Work by the Argonne group has shown that for Ar/H2 /CH4 (1% CH4) discharges the transition from a discharge containing no argon (hydrogen dominant) to one containing no hydrogen (argon dom-
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inant) is accompanied by fundamental changes in the plasma chemistry, growth process, growth rate, and deposited film properties. However, throughout the entire variation of hydrogen-argon compositions with a small percentage of methane (⬃1%), diamond is grown. The plasma chemistry is observed to change from a plasma dominated by CH3 and H radicals at high hydrogen concentrations to a plasma having a high C2 concentration at high argon concentrations. Corresponding to the discharge chemistry changes, the diamond growth scenario for 70–99% hydrogen is dominated by the CH3 growth species, for the 20–70% hydrogen region the growth rate is dominated by the C2H2 species, and for the low hydrogen content region (less than 20%) the growth is dominated by the C2 species [59]. The C2 species is postulated as both a diamond growth species and a diamond nucleation species. The use of a predominantly argon discharge with small amounts of CH4 (⬃1%) and small amounts of hydrogen (0 to a few percent) permits the deposition of very smooth diamond films with randomly oriented 3- to 10-nm crystallites of diamond. Such diamond is referred to as nanocrystalline diamond.
VI.
INTERNAL VARIABLES
Several researchers have studied the internal variables in microwave plasmaassisted CVD diamond deposition reactors. These studies have quantified some of the relationships of input to internal variables. Some of the internal variable quantities measured have been species concentrations, gas temperature, electron density, electron temperature, and microwave power absorption density versus various input parameters. In this section these internal variable studies are reviewed with an emphasis on the studies that quantified the internal variable values. In particular, optical emission studies that provide only relative signal levels or relative concentration levels are not considered in this review. A.
Species Concentrations
Mitomo and coworkers [60] used in situ Fourier transform infrared (FTIR) spectroscopy to measure the molecular species concentrations in the microwave discharge region. In their study they used a number of different input gas mixtures including CH4 ⫹ H2 , C2H2 ⫹ H2 , C2H4 ⫹ H2 , CH3OH ⫹ H2, C2H5OH ⫹ H2 , CO ⫹ H2 , CCl4 ⫹ H2 , and CH4 ⫹ O2 ⫹ H2 for the discharge. The reactor was a tubular unit, shown in Fig. 5, operated at 2.45 GHz with a 4-cm-diameter tube, 30 torr pressure, and 300 W microwave power. The carbon-containing species observed using FTIR included CH4 , C2H2 , C2H4,
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and CO. The CH4 and C2H2 species were observed under all the conditions studied, while C2H4 was observed only when higher concentrations of carbon input gases were used. Lastly, CO was produced and measured when oxygen was present in one of the input molecular gas species. The species CH3OH, C2H5OH, and CCl4 were not seen in the discharge, indicating that plasma conditions were effective at dissociating these molecules. As shown in Fig. 37, when the input gas contained no oxygen the dominant species for low percentages of input carbon gas flow (i.e., 0.5% or less) was methane and the dominant species for percentages above 0.5% was acetylene. This was true regardless of the molecular form of the input gas. Another observation was that the addition of oxygen to the discharge reduced the concentrations of CH4 and C2H2 in the discharge as compared with the levels when no oxygen was present. However, the CO species concentration increased as oxygen was added. Hence, the discharge conditions in microwave CVD
Figure 37 Concentration of CH4 , C2H2 , and C2H4 in a microwave discharge region as measured by FTIR absorption spectroscopy versus percent methane in the feed gas at an operating pressure of 30 torr with no oxygen in the feed gas. (From Ref. 60.)
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reactors favor the formation of CO; i.e., the addition of oxygen results in bonding of carbon with oxygen, forming CO. A mass spectrometric study by Weimer et al. [44] found similar results. They used a differential pumped quadrapole mass spectrometer to sample the exhaust gas from a microwave tubular reactor being operated for the deposition of diamond. The reactor had a 1.25-cm-diameter tube operating at a pressure of 21 torr, a microwave power of 170 W, and a frequency of 2.45 GHz. The input gas flow consisted of hydrogen, oxygen, and various carbon-containing molecules including CH4 , C2H2 , C2H4 , and C2H6 . The species detected in the exhaust gas included CH4 , C2H2 , CO, and H2O. Figure 38 shows plots of the various detected species as a function of the oxygen percentage in the input gas flow (0–3%). Once again, the dominant carbon species at low oxygen input was either CH4 or C2H2 . At oxygen flow rates of 1% or greater the dominant species was CO and the concentrations of methane and acetylene dropped as the oxygen flow increased. Using a molecular beam mass spectrometer that is capable of measuring radical and stable species in the discharge, Hsu [61] studied diamond deposition in an ASTeX microwave reactor (see Fig. 7). The experimental structure was constructed so that the sampling point of the discharge was a hole in the center of the substrate holder. The deposition conditions used during the mass spectrometer measurements included a substrate temperature of 1073 K, a pressure of 20 torr, and an input microwave power of 800 W. The species detected by the spectrometer included H2 , H, CH3 , CH4 , and C2H2 . The results, which are summarized in Fig. 39, show measured concentrations versus the methane input flow percentage from 0.2 to 3%. The two radicals measured were atomic hydrogen and the methyl radical, CH3 . The dominant carbon species at low methane input flow was methane, whereas at methane flow rates of 0.5% and greater the dominant species was acetylene. The measured hydrogen mole fraction was about 10⫺3. This, of course, is the measured value at the surface of the substrate holder where the discharge is sampled by the spectrometer. It was further estimated by Hsu that the H-atom mole fraction in the center of the discharge located above the substrate holder would be 8%. He also found that the hydrocarbon chemistry in the microwave system was largely unchanged from that found in thermal systems (i.e., hot-filament systems). This was due to the carbonhydrocarbon chemistry being driven by the gas temperature and the atomic hydrogen concentration. The primary effect of the plasma was suggested to be the dissociation of the hydrogen molecules to atomic hydrogen. Erickson et al. [62] measured in a unique microwave reactor configuration other radical species concentrations in a microwave discharge reactor. The measurement technique used was absorption spectroscopy. The species measured included CH3 and CH at a pressure of 20 torr in the discharge.
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Figure 38 Diamond deposition reactor exhaust gas composition versus oxygen percentage in the feed gas for input gas flows of 97 sccm hydrogen and (a) 2 sccm CH4 , (b) 1 sccm C2H6 , (c) 1 sccm C2H4 , and (d) 1 sccm C2H2 . (From Ref. 44.)
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Figure 39 Mole fraction of chemical species measured in a microwave plasma CVD system versus variation in the methane fraction in the feed gas. Measurements done using a molecular beam mass spectrometer. (From Ref. 61.)
Based on the measured values of CH3 and CH at 20 torr, the atomic hydrogen mole fraction was estimated to be 0.8% or less. Wouters and coworkers [63] looked at hydrogen discharges created in a 1.6-cm-diameter tubular microwave reactor operating in the pressure range from 2 to 40 torr. They used the two-photon laser-induced fluorescence (LIF) method to measure the atomic H density when this discharge was operated in a pulsed mode with a 40% duty cycle at 100 Hz using 200 W of power. The results obtained for the absolute atomic hydrogen density and mole fraction are shown in Fig. 40. They reported gas temperatures in the discharge region were 1600 K at 40 torr and 1200 K at 2 torr as measured by H-atom Doppler broadening. The atomic hydrogen mole fraction can be seen to vary from 8% at 2 torr to 2% at 40 torr. Atomic hydrogen mole fraction in an LIMHP-France bell jar microwave CVD reactor (shown in Fig. 10) was measured by Gicquel and coworkers [26,64]. They used the method of actinometry, where the emission signals are properly calibrated so that absolute densities can be obtained. Actinometry is a technique that looks at the strength of optical emission
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Figure 40 63.)
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Atomic hydrogen percentage and density versus pressure. (From Ref.
signals from the discharge. Specifically, the emission from atomic hydrogen and the emission from a gas added at a low concentration (single-atom gas such as argon) are both measured from the same discharge. If the electron excitation energy and cross section of excitation of the two emitting species are similar, then the density of the atomic hydrogen species can be determined because the density of the argon species is known from the input flow rate. Further, if the signal intensities are properly calibrated, the absolute densities can be determined. In this case the results obtained for the atomic hydrogen concentration are shown in Fig. 41 [26,64]. The system used for this measurement was a microwave resonant reactor with a 5-cmdiameter substrate and a 10-cm-diameter discharge region. The power density variation (6–30 W/cm3) was obtained by varying the pressure and the input power to obtain the same plasma volume as the pressure was increased from 18 to 100 torr. The power density and pressure are strongly coupled with pressure increases producing power density increases. As shown in Fig. 41a, the atomic hydrogen mole fraction increases from about 2 to 20% as the pressure increases. Figure 41b displays the variation of the atomic mole fraction of hydrogen as the methane concentration is increased from 0 to over 5%. Only a slight decrease is observed at the low power density of 9 W/cm3. At the higher power density of 22.5 W/cm3, no significant change in the atomic hydrogen concentration is observed as the methane concentration increases from less than 1% to 6%.
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Figure 41 Atomic hydrogen mole fraction versus the microwave power density in part (a) and versus methane concentration in part (b). (From Ref. 26.)
B.
Neutral Gas Temperature
The neutral gas temperature of the discharge in microwave plasma-assisted diamond deposition reactors has been measured in a number of ways including Doppler broadening of optical emission signals, Doppler broadening of LIF signals, coherent anti-Stokes Raman spectroscopy (CARS), and rotational temperature of optically emitting molecules. One of the most systematic studies of the gas temperature in a diamond deposition reactor has been done by Gicquel and coworkers [65,66]. They used all four of the techniques just listed to measure the gas temperature in an LIMHP-France bell jar microwave discharge reactor (Fig. 10) operating at 2.45 GHz. This reactor had a discharge diameter of 10 cm and a substrate diameter of 5 cm.
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The results found by measuring the line-of-sight Doppler broadened emission of atomic hydrogen are shown in Fig. 42. The gas temperature ranges from 1600 K at the low power density (18 torr) to 3500 K at the highest power densities studied, which correspond to a pressure of 105 torr. Measurements performed using laser spectroscopy techniques are shown in Fig. 43. The two techniques utilized were CARS to measure molecular hydrogen and two-photon LIF to measure atomic hydrogen. The CARS technique measures the rotational temperature of molecular hydrogen. The LIF technique measures the Doppler broadening due to atomic hydrogen translational motion. Both of these techniques can be performed with good spatial resolution. For the data shown in Fig. 43 the location of the measurement was in the center of the discharge at a distance of 20 to 25 mm above the substrate. The effect of methane addition on the gas temperature was also studied as shown in Fig. 44. The results obtained by the line-of-sight optical
Figure 42 Atomic hydrogen translation temperature versus microwave power density. The temperature was measured using Doppler broadening of hydrogen OES signals. (Adapted from Refs. 65, 66.)
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Figure 43 Gas temperature versus average microwave power density. Gas temperatures measured include atomic hydrogen temperature by two-photon LIF and hydrogen rotational temperature by CARS. (Adapted from Refs. 65, 66.)
emission measurements of the rotation spectrum of molecular hydrogen are shown in Fig. 45. The temperatures in this case range from 1300 K at 9 W/ cm3 to 1900 K at 22 W/cm3. These rotational temperature values are systematically lower than those found by the other techniques. Similar rotational temperatures versus pressure were also measured by Grotjohn et al. [67] for the MCPR reactor system of Fig. 6b. One variation of this reactor has a 5cm discharge chamber diameter and a 3-cm-diameter substrate holder. These results are shown versus pressure in Fig. 46. Some understanding of the difference in the temperature measurements can be obtained by considering the location and species being measured. The Doppler broadening measurement of atomic hydrogen is weighted toward the region that has the highest atomic hydrogen concentrations. That region is the center of the plasma, which has the highest gas temperature. Conversely, the molecular hydrogen, which is the species being observed
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Figure 44 Atomic hydrogen temperature (measured by LIF) and molecular hydrogen rotational temperature (measured by CARS) versus methane percentage in the feed gas. (Adapted from Refs. 65, 66.)
for the OES (optical emission spectroscopy) rotational temperature measurements, will have a higher density near the edges of the discharge where the gas temperature is lower. Hence, the line-of-sight rotational temperature measurements are weighted more toward the edges of the plasma discharge, which are expected to be cooler. C.
Electron Density and Temperature
The electron density in the discharges operated under conditions that deposit diamond have been measured. Grotjohn and coworkers [67] used a millimeter wave Fabry-Perot resonator operating at 30 GHz to measure the electron density in a 2.45-GHz microwave reactor operating in the pressure range from 10–60 torr. The results of the measurements are shown in Fig. 47 versus pressure. The density range is 1–5 ⫻ 1011 cm⫺3. In another investigation using a tubular reactor with a diameter of 2 cm, Tahara et al. [68] studied methane-hydrogen discharges in the pressure range from 0.1 to
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Figure 45 Molecular hydrogen rotational temperature versus average microwave power density. Temperature was measured using light-of-sight OES. (Adapted from Refs. 65, 66.)
5 torr. The tube was placed through the center of a 2.45-GHz tunable resonant cavity. The methane flow rate in the input gas flow was varied from 1 to 10%. As shown in Fig. 48, the plasma density, which was measured with a Langmuir probe, stayed approximately constant at 1 ⫻ 1011 cm⫺3 across the pressure range studied. The microwave power utilized was 350 W. The last internal variable quantity is the electron temperature, which was measured in the pressure range 0.1 to 5 torr by Tahara et al. [68] using the system described in the preceding paragraph. The electron temperature values varied from slightly above 3.5 eV at the lower pressure of 0.1 torr to about 2eV at the higher pressure of 5 torr. D.
Microwave Power Density
Another important internal variable is the power density of the plasma discharge. Power density is defined by knowing the absorbed power of the
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Figure 46 Molecular hydrogen rotational temperature versus pressure for a 5-cmdiameter discharge system. (From Ref. 67.)
discharge and dividing by the volume of the discharge. The volume is typically estimated from the region of most intense light emission by the discharge. An example plot of the power density versus pressure measured for the MCPR reactor of Fig. 6b operated with a 3- to 4-cm diameter discharge and 3.2-cm-diameter substrate holder is shown in Fig. 49 [67]. Even though the pressure is a primary determining factor for the power density, other factors such as substrate size and total plasma volume are also determining factors. For example, the power density of the microwave system studied here does not match exactly the power density of other microwave diamond CVD systems with different diameters of substrate holders. In work by Kuo and Asmussen [8], the power density found for a microwave reactor with a 7.5-cm-diameter holder at 60 torr was 15 W/cm3 (see Fig. 17). In another system [2,30,65,66], the power density for a microwave system with a 5cm-diameter substrate holder operating at 60 torr was 22.5 W/cm3. So at the 60 torr pressure and 2.45 GHz excitation frequency, microwave systems with substrate holder diameters of 3.2, 5, and 7.5 cm have power densities of 32, 22.5, and 15 W/cm3, respectively. This relationship indicates that larger diameter plasmas, which are needed to cover larger substrates, operate with a lower power density than smaller diameter plasmas.
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Figure 47 Electron density versus pressure for a microwave plasma-assisted diamond deposition discharge. (From Ref. 67.)
E.
Microwave Electric Field and Mode
The last quantity to be considered in this section is the measurement of the impressed microwave electric field strength and the spatial variation of the electromagnetic mode in diamond deposition reactors. The specific system considered here is the microwave cavity plasma reactor (MCPR) as shown earlier in Fig. 6b. The microwave electric field has been measured in the MCPR system by inserting a microcoaxial antenna probe as shown in Fig. 50 through holes drilled in a pattern along the applicator cavity walls, i.e., along the vertical or z direction and around the circumference. The microcoaxial antenna probe samples the electric field component normal to the wall, which in this case is the radial component. The antenna probe is inserted through the applicator cavity holes a small distance (⬃a few millimeters) so that only a small fraction of microwave power going into the applicator cavity is coupled into the sampling antenna probe (typically
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Figure 48 Electron temperature and plasma density versus pressure for a cavity resonance microwave hydrogen-methane discharge. (From Ref. 68.)
⬃10⫺6). This small fraction for the amount of power sampled by the antenna probe ensures that the probe does not perturb the field in the applicator. An example plot of the measured electric field strength versus cavity vertical position (z direction) is shown in Fig. 51 [69,70]. The symbols show the experimentally measured microwave electric field values in a discharge loaded applicator cavity. The dotted line is the expected strength of the radial field variation versus vertical position along the applicator cavity for the TM013 mode. A sketch of the ideal TM013 mode is shown in Fig. 52. Note in this figure that the electric field has three maximums along the longitudinal direction, which is also seen in the experimental data of Fig. 51. Although not shown in Fig. 51, when the microwave electric field is sampled around the circumference of the applicator cavity wall, the electric field strength is unchanged in agreement with the MPCR operating in a TM013 electromagnetic mode.
VII.
MICROWAVE DIAMOND CVD REACTOR MODELING
Microwave plasma-assisted CVD reactors for diamond deposition have been modeled in a number of investigations. The modeling of microwave CVD reactors needs to consider a number of factors including:
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Figure 49 Microwave power density versus pressure for 5-cm-diameter hydrogenmethane discharge used for diamond deposition. (From Ref. 67.)
1. 2. 3.
Electron gas–initiated reactions in the discharges Neutral chemistry (including thermally activated) reactions in the discharges Microwave field excitation, propagation/resonance, and absorption including the spatial patterns of microwave fields and microwave power absorption
First, a few preliminary comments regarding H2 /CH4 discharge and microwave discharge operation in the pressure range from 1 to hundreds of torr are in order. The dissociation reactions, which are so important to diamond deposition (i.e., dissociation of H2 and CH4), can be driven by either electron dissociation or thermal dissociation. The general trend as shown in Fig. 53 for all molecular gases is that at low pressures, the electron temperature is high and the gas temperature is low so electron dissociation dominates. Conversely, at high pressures, the electron temperature is reduced
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Figure 50 Details of the microwave electric field measurement probe and a cutaway view of a microwave cavity plasma reactor.
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Figure 51 Microwave electric field strength (radial direction) versus vertical position for a microwave cavity plasma reactor as shown in Fig. 50.
Figure 52
Sketch of TM013 electromagnetic resonant mode.
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Figure 53 Electron, ion, and neutral temperatures versus pressure for nonequilibrium (nonisothermal) and thermal-like plasmas. (From Ref. 10.)
and the gas temperature increases, resulting in increased thermal dissociation. The exact pressure range over which this transition occurs is dependent on many factors including molecular species type, other constituents in the plasma, and the plasma volume. One of the objectives in this section is to establish an understanding of this transition in the diamond deposition process. A.
Modeling Studies
Several researchers have studied diamond deposition in microwave plasmaassisted CVD systems. The key attributes that distinguish the various models are: 1. 2.
Spatial dimension of the solution [zero (one point), one, or two dimensional] Discharge chemistry model (hydrogen, hydrogen-methane, H-CO)
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Microwave power absorption profile (uniform, assumed profile, self-consistently solved profile)
Koemtzopoulos et al. [71] in 1993 reported on a model for microwave discharges that looked at hydrogen dissociation. The reaction rates were solved for by determining the electron energy distribution function (EEDF) using the Boltzmann equation. The model considered electron dissociation of the hydrogen as the primary dissociation mechanism. The model solved the electron density balance equation, the atomic hydrogen balance equation, and the power balance equation at pressures ranging from 15 to 100 torr. The geometry considered was a tubular reactor with the electric field and species concentrations assumed uniform throughout the plasma volume. The model showed that the electron distribution function is non-Maxwellian with the tail of the distribution function depleted by inelastic collisions. Further, it was demonstrated that even when 1% CH4 was added to the hydrogen discharge of EEDF did not change appreciably. This work also noted that the dissociation of methane by abstraction of atomic hydrogen (CH4 ⫹ H → CH3 ⫹ H2) is faster than electron dissociation at pressures as low as 20 torr with a gas temperature of 600 K and an atomic hydrogen fraction of 0.08. Hyman et al. [72] in 1994 developed a one-point numerical model of microwave plasma CVD diamond deposition reactors. The model solved the Boltzmann equation to determine the electron energy distribution function. A dominant energy loss mechanism for the electron gas was the vibrational excitation of the hydrogen molecules. The gas heating then occurred predominantly through vibrational-translational energy exchange. This model considered discharges consisting of carbon-hydrogen-oxygen mixtures. In the model the point of calculation was selected as the center of the plasma ball located above the substrate holder. Diffusion and heat conduction effects were accounted for by two boundary points fixed at the substrate condition and the wall/ambient condition. Their simulation example was given at 40 torr pressure, 30 W/cm3 microwave power density, and an input gas mixture of 98.75% H2 , 1% CH4 , and 0.25% O2 . Typical values of the plasma conditions simulated at time scales approximating steady state (after 4 msec of reaction time) are an electron temperature of 2.5 eV and a gas temperature of ⬃3000 K. In their model the initial gas mixture of H2 , CH4 , and O2 reacts to a mixture consisting of H2 , C2H2 , and CO (stable species) and H, C, CH3, and C2H (radical species). Hence the methane and oxygen molecules in the discharge are converted predominantly to C2H2 and CO molecules. Work by Tan and Grotjohn [69,73] on microwave plasma reactors for diamond deposition looked at the microwave fields and the absorption of the microwave energy by the discharge. They obtained full-wave two-di-
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mensional solutions of Maxwell’s equations to solve for the electromagnetic fields using the FDTD (finite-difference time-domain) method. The work simulated the resonant fields in a resonant cavity microwave reactor. Experimental measurements and simulations of the 2.45-GHz microwave fields showed close agreement. This model was limited to hydrogen discharges in the reactor. The reaction kinetics included electron dissociation and ionization processes of hydrogen, but thermal processes were either left untreated or treated in a very simple manner. The gas temperature in their model was set on the basis of experimental gas temperature measurements. This modeling work solved for the electromagnetic fields and the power absorption of the electron gas in the microwave reactor. Extensive modeling work on hydrogen and hydrogen-methane discharges has been done by Hassouni et al. [74–80] at LIMHP-CNRS in France. They developed one-dimensional discharge models for moderatepressure H2 and H2 /CH4 plasma discharges. They also developed a twodimensional model of H2 discharges using an assumed microwave power absorption profile. Their modeling work considered the electron energy, H2 vibrational energy, and total energy. The models use Boltzmann equation solutions to determine the EEDF, which is important for determining electron reaction rates because the EEDF is non-Maxwellian. The model also included gas heating in the discharge region. Extensive comparisons with good agreement are made of the simulated gas temperature and the experimentally measured gas temperature, which was obtained by CARS and Doppler broadening resolved OES. The hydrogen discharge considered in their modeling work included either 7 or 9 species and the H2 /CH4 model included an additional 12 neutral carbon-containing species and 9 ionic carbon species. The carbon chemistry for the pressure range considered (18– 100 torr) has been shown to be primarily governed by thermal chemistry with the gas temperature and the atomic hydrogen concentrations being the primary neutral carbon chemistry driving factors. The work by Tan and Grotjohn [69,73] on modeling the microwave fields in microwave diamond deposition reactors was combined with the work on hydrogen discharges by Hassouni et al. [74–80] to produce a twodimensional self-consistent microwave field and plasma discharge simulation for a moderate pressure (18–80 torr) hydrogen discharge reactor [81]. This model produced a self-consistent solution of the microwave fields in the plasma reactor, the absorption of the microwave energy by the discharge electron gas, and the hydrogen plasma discharge. The plasma model used was a three energy mode (gas, molecular vibration, and electron) with nine ⫹ ⫺ ⫺ species [H2 , H, H (n = 2), H (n = 3), H⫹, H⫹ 2 , H3 , H , and e ] model that accounted for the non-Maxwellian electron distribution function and had 35 reactions. Good agreement was obtained between the model results and the
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experimental data for gas temperatures and atomic hydrogen concentrations. The model was also able to predict plasma discharge (ball) size, shape, and location in the discharge chamber with good accuracy. The specific reactor simulated, as shown in Fig. 10, had a 25-cm-diameter microwave resonant cavity structure operating at 2.45 GHz with a silica bell jar of 10 cm diameter. The substrate holder inside the bell jar had a diameter of 5 cm as shown in Fig. 10. Example two-dimensional plots of the gas temperature, atomic hydrogen mole fraction, absorbed microwave power density, electron temperature, and electron density are shown in Fig. 54. Numerical simulations of microwave plasma reactors for diamond CVD have also been done by Funer et al. [82,83]. In their study microwave fields are modeled in two dimensions (r and z) using a finite integration technique to solve Maxwell’s equations for the reactor structure. The plasma model used was a cold plasma dielectric function description of the discharge where the density of the discharge was calculated on the basis of the local electric field and a fitting parameter ␥ that is adjusted to have the modeling results agree with experimental observations. Their work looked at using the numerical model to design an optimum reactor. The various geometric dimensions of the reactor are described by parameters P1 , P2 , . . . , Pn . An optimization is done to determine the set of parameter values that maximize the ratio of the electric field (兩E兩2) within the plasma region to the electric field strength (兩E兩2) in the region surrounding the plasma region. The optimum parameter set was determined using the method of ‘‘direction of steepest gradient’’ to find the optimum solution. The result of this study was the elliptical reactor described earlier in Sec. IV, Fig. 11 [16]. B.
Internal Variables: Modeling Results
From the various modeling efforts, a number of important understandings regarding microwave-assisted diamond CVD reactors can be extracted. The first is that the chemistry of neutral carbon species (e.g., dissociation of CH4 to CH3) is essentially driven by gas temperature and H-atom concentration. Hassouni et al. [74] demonstrated using a one-dimensional H2 /CH4 plasma model that for the pressures they studied (18 torr and above) the carbon chemistry is driven by thermal and atomic hydrogen [H] processes and not by electron excitation or dissociation of carbon species. Another understanding that can be drawn from the various modeling results is the relative importance of thermal dissociation and electron collisional dissociation of hydrogen. First, Fig. 55 shows the atomic hydrogen concentration under equilibrium conditions due to just thermal dissociation at temperatures from 1600 to 2800 K (dashed line) for the two pressures of 20 and 100 torr. The equilibrium molar fraction for a pressure of 20 torr
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Figure 54 Spatial distribution (versus r and z) of discharge characteristics at 2500 Pa and 600 W. (a) Gas temperature (K), (b) atomic hydrogen model fraction, (c) absorbed microwave power density (W/cm3), (d) electron temperature (K), and (e) electron density (⫻1011 cm⫺3). (From Ref. 81.)
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Figure 55 Hydrogen dissociation (molar percentage) versus gas temperature for hydrogen in equilibrium conditions (dashed line) and hydrogen in a plasma discharge used for diamond deposition (solid line).
goes from 0.038% at 1600 K to 1.1% at 2000 K to 41% at 2800 K. Data are also shown in Fig. 55 for a pressure of 100 torr. These quantities are taken from the work by Argoitia et al. [84]. In the two-dimensional simulations of a 2.45-GHz microwave plasma-assisted CVD reactor with a discharge size of approximately 5 cm diameter by Hassouni et al. [81] the simulation results yielded an average hydrogen molar fraction and gas temperature as a function of pressure as shown in Fig. 56. Note that as the pressure increases the gas temperature and the hydrogen molar fraction also both increase. If these gas temperature and hydrogen molar fraction values are now also plotted (Fig. 55), the solid line labeled 2.45 GHz is obtained. In Fig. 55, the lowest pressure of 18 torr shows that the hydrogen dissociation in the microwave CVD system is substantially greater that that expected from the gas temperature. Hence, at this lower pressure electrondriven dissociation of the discharge is an important process. At the higher pressures, the hydrogen molar fraction is approximately the same as or less than the thermal dissociation value based on the discharge gas temperature. Hence, at 40 torr and above the thermal dissociation of hydrogen becomes a dominant process. The modeled/measured value of the hydrogen dissociation is less than the equilibrium value because, in the reactor, the discharge chamber walls and cooler gas regions outside the active discharge volume provide sinks for atomic hydrogen via surface recombination and volume recombination. The influence of reactor scaling on the atomic hydrogen molar fraction is also interesting to consider. A larger diameter microwave diamond dep-
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Figure 56 (a) Plasma gas temperature versus pressure in a hydrogen-dominated diamond deposition discharge. (b) Atomic hydrogen percentage versus pressure in a hydrogen-dominated diamond deposition discharge.
osition reactor operating at 915 MHz with a substrate diameter of 15 cm was also simulated using the model in Ref. 81. The simulated data for the larger system are shown in Fig. 55 as a solid line labeled 915 MHz. It can be seen that at the higher gas temperatures of 2400 K and above the larger diameter 915 MHz system has an atomic hydrogen molar percentage that is closer to the thermal equilibrium values (dashed lines) than the smaller 5cm-diameter discharge 2.45-GHz system. Another important factor affecting the %H and the gas temperature at a selected pressure is the input microwave power. Table 2 shows the gas temperature versus input microwave power at a pressure of 18 torr in a
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Table 2 Simulated Plasma Discharge Quantities as a Function of Input Microwave Power at a Pressure of 2500 Pa Microwave input power (W) Discharge parameters Vp (cm3) Tg peak value (K) Tg average value (K) Tg peak location (K) Te peak value (K) Te average value (K) Te peak location (cm) MWPD peak value (W/cm3) MWPD average value (W/cm3) MWPD peak location (cm) ne peak value (cm⫺3) ne average value (cm⫺3) ne peak location (cm) xH peak value (%) xH average value (%) xH peak location (cm)
300
500
700
800
900
19 1,730 1,500 1.3 19,840 16,030 0.3 18.1 8.7 0.3 3.3 ⫻ 1011 2.0 ⫻ 1011 0.9 0.65 0.4 1.1
33.6 1,920 1,620 1.5 19,920 16,000 0.3 19.2 8.5 0.3 3.4 ⫻ 1011 2.1 ⫻ 1011 1.1 0.81 0.52 1.5
49 2,020 1,640 1.7 19,650 15,930 0.3 18.4 8.0 0.3 3.3 ⫻ 1011 2.1 ⫻ 1011 1.3 0.9 0.55 1.5
61 2,080 1,710 1.9 19,570 15,730 0.3 17.5 7.4 0.3 3.3 ⫻ 1011 2.1 ⫻ 1011 1.3 0.93 0.60 1.7
74.5 2,065 1,640 1.9 19,560 15,320 0.3 17.2 6.4 0.3 3.1 ⫻ 1011 1.8 ⫻ 1011 1.1 0.9 0.52 1.7
Source: Ref. 81.
hydrogen discharge [81]. The primary factors affecting the gas temperature and percent hydrogen concentration once a particular pressure is selected are the plasma volume and the plasma boundary temperature (substrate temperature and wall temperatures). In general, the larger the plasma volume (as determined by the discharge chamber size and the input microwave power), the higher will be the gas temperature. The variations of other internal variables versus microwave input power are also given in Table 2. A companion table showing the variation of the internal variables versus pressure for hydrogen discharges of similar volumes is Table 3. Tables 2 and 3 provide a picture of the internal variable values and their variation with the pressure and microwave power (two input variables). One final utility of the models is in the design and operation of a reactor to get a uniform deposition across large areas. It has been noted by researchers that the uniformity of the deposition is controllable by the adjustment of the substrate position, microwave tuning circuit, pressure, and amplitude of the input microwave power. For example, work described by Bachmann [10] showed that films can be grown in the same reactor that vary from thickest in the center of a wafer area, to relatively uniform, to
294 Table 3
Grotjohn and Asmussen Simulated Plasma Discharge Quantities as a Function of Pressure/Power Pressure (hPa), microwave input power (W)
Discharge parameters 3
Vp (cm )* Tg peak value (K) Tg average value (K) Tg peak location (cm) Te peak value (K) Te average value (K) Te peak location (cm) MWPD peak value (W/cm3) MWPD average value (W/cm3) MWPDav experiment (W/cm3) MWPD peak location (cm) ne peak value (cm⫺3) ne average value (cm⫺3) ne peak location (cm) xH peak value (%) xH average value (%) xH peak location (cm)
25–600
52–1000
84–1500
110–2000
44.1 1,990 1,640 1.7 19,680 15,900 0.3 18.3 8.0 9.0 0.3 3.4 ⫻ 1011 2.1 ⫻ 1011 1.1 0.87 0.55 1.5
40.7 2,600 2,100 1.7 cm 17,220 14,600 0.3 40.1 16.0 15.0 0.3 5.1 ⫻ 1011 3.1 ⫻ 1011 1.1 1.19 0.78 1.5
39.8 3,100 2,540 1.7 cm 15,680 13,600 0.3 67.8 25.0 22.5 0.3 7.4 ⫻ 1011 4.3 ⫻ 1011 1.1 2.91 1.66 1.7
45 3,200 2,760 1.7 cm 13,990 12,400 0.3 79.1 30. 30. 0.3 9.1 ⫻ 1011 5.3 ⫻ 1011 1.1 6.98 3.8 1.9
*Vp, plasma volume. Source: Ref. 81.
thickest at the edge of the deposition area. Such a dependence of the uniformity was also seen in the discharge itself in Ref. 81, where simulations with the substrate holder at three difference vertical positions yielded the results of Fig. 57. Specifically, variation of the substrate height produced changes in the gas temperature and in the shape of the plasma discharge region above the substrate. Raising the substrate holder is shown in Fig. 57 to flatten the plasma so that it is more uniform across the diameter of the substrate.
VIII.
CONCLUSION
Microwave plasma-assisted CVD machines create and control the chemical environment and substrate conditions for diamond growth. The appropriate chemically active environment supplies neutral fluxes, such as CH3 and C2H2 , that provide the carbon for the growth process and also supplies a high-density atomic hydrogen flux so that deposition occurs in the diamond
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Figure 57 Spatial distribution of gas temperature (K) at three substrate positions. (a) Lowered 1 cm, (b) nominal position, and (c) raised 1 cm. (From Ref. 81.)
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phase. The deposition machine must also control the substrate temperature to achieve the desired diamond film uniformity and properties. Since the early 1980s, when Kamo et al. [6] created the tubular reactor, substantial progress and major innovations have been made in the design and development of microwave plasma-assisted CVD machines. A wide range of machine types and sizes have been developed and have demonstrated successful diamond deposition. Using current technology and know-how, microwave plasma-assisted CVD is achieved over a wide experimental regime, including pressures ranging from 10 m torr to 1 atmosphere, microwave powers ranging from 100 W to 60 kW, and substrate diameters ranging from 2.5 to 30 cm. Today, microwave CVD systems can deposit diamond films uniformly across substrates as large as 30 cm in diameter and at rates of nearly 1 g/hr. By carefully adjusting the chemistry and controlling impurities, diamond films with the desired quality and morphology can also be synthesized. Considerable engineering has gone into the design of these systems, and reliable, automated operation is currently available for commercial applications. The approach utilized in this chapter was to view these diamond machines in terms of the relationships between the input variables, internal variables, and output variables. The relationships between these experimental variables are developed from both experimental studies reported over the past 15 years in the literature and modeling investigations of specific diamond deposition machines and their associated diamond deposition processes. Much has been learned from experimental and modeling investigations about the relationships between the input variables, such as pressure, gas flow rate and composition, and input microwave power, and the internal variables, i.e., chemistry, of the discharge. This understanding has come both from careful experimental measurements of the internal variables including gas temperatures, plasma density, radical densities, and microwave power densities and from careful modeling studies that can now solve for the plasma discharge kinetics and the microwave fields self-consistently in two dimensions. These studies, which were reviewed in this chapter, provide a very useful picture for the understanding of the engineering and scientific issues associated with the design of microwave plasma-assisted CVD machines and the diamond deposition process itself. Microwave plasma reactors are often used for fundamental scientific experimental investigations because microwave plasma-assisted machines operate without electrodes or filaments in contact with the discharge. Hence, electrode, filament, and other contamination issues can be minimized. For example, microwave plasma reactors have been extensively used to investigate the influence of impurities (even at levels as low as 10 ppm) on the deposition process and film properties. One result of these investigations has
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been the design and development of very reliable, low-maintenance microwave plasma-assisted CVD machines that have very high vacuum system capabilities as well as the ability to control impurity gases to levels below a few ppm. Thus applications that require high-quality diamond synthesis, such as diamond window synthesis as well as synthesis of electronic quality films, are now usually carried out using microwave plasma reactors. However, even with the advances and successes over the past 15 years in the areas of machine design and process understanding, many interesting design and process understanding issues remain. Some of these include understanding the influence of various impurities or additives on the deposition process, e.g., the influence of small amounts of nitrogen, boron, and oxygen across the range of deposition conditions. Other issues center around the design of deposition machines that permit ever more uniform diamond deposition across large areas while minimizing unwanted contamination or impurities from the machine itself, i.e., impurities from wall erosion, substrate erosion, and fused silica bell jar or window erosion found at the high temperatures and in the highly active chemical environments of the deposition process. REFERENCES 1. 2.
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8 Combustion Synthesis of Diamond Colin A. Wolden Colorado School of Mines, Golden, Colorado
I.
INTRODUCTION
Combustion synthesis is one of the competing chemical vapor deposition (CVD) technologies for diamond film growth. It was first discovered in 1988 by Hirose and Kondo [1], who realized that acetylene flames produce copious amounts of atomic hydrogen and hydrocarbon radicals that are required for diamond growth. The discovery was subsequently confirmed at the Naval Research Laboratory [2] and has become the subject of intense study. It has been argued that combustion synthesis is the most flexible of the CVD alternatives because of its scalable nature, minimal utility requirements, and significantly reduced capital costs relative to plasma-aided processes [3]. In this chapter the development of combustion CVD will be reviewed from its inception using conventional welding torches to its present implementation with flat flame burners. The key experimental parameters and their impact on deposition are discussed in detail. The chemistry of combustion diamond CVD is discussed, highlighting studies that employed in situ diagnostics and detailed reactor modeling. Finally, combustion synthesis is contrasted with other CVD techniques in terms of both processing and economics.
II.
INVENTION: THE WELDING TORCH
A.
Description of System
Figure 1 shows the structure of an atmospheric oxyacetylene torch in its unperturbed state and in the configuration used for diamond deposition. 303
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Figure 1 Schematic diagrams of (a) the flame structure of an oxyacetylene torch and (b) the flame in a diamond deposition configuration.
There are three distinct regions in an acetylene flame: (1) the inner flame, (2) the acetylene feather, and (3) an outer diffusion flame. It is in the feather region that the substrate is placed for diamond growth. The most important parameter in combustion synthesis is the oxygen-to-acetylene ratio, defined here as R = O2 :C2H2 . At values of R near 1.0, a neutral flame is achieved, which is defined as the condition where the feather region just disappears because all the acetylene is consumed in the primary flame. The diamond growth regimes as a function of composition are outlined in Fig. 2. The highest quality diamond is obtained in slightly rich flames, R = 0.85–1.0 [4]. The value of R at which a neutral flame occurs is dependent on both burner design and total flow rate. Thus, for a particular experimental configuration the composition at the neutral flame condition may be shifted slightly, but the principles described in Fig. 2 still hold. The welding torch arrangement of Hirose has been used in numerous subsequent studies [5–13]. Polycrystalline diamond has been grown at rates up to 200 m/hr [5] and thick (150 m) homoepitaxial layers have been deposited [6]. Atmospheric torches have successfully produced large indi-
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Figure 2 Carbon deposition regions observed using atmospheric acetylene torches as a function of gas composition. (From Ref. 4.)
vidual crystals that approach 1 mm in diameter [7,8]. The drawback of the welding torch arrangement is that the deposited material is characterized by radial inhomogeneity [9] and often contains significant amounts of nondiamond carbon [10]. Because of these limitations, welding torches are being replaced by flat flame burners, which are more appropriate for large area deposition. Nevertheless, the wealth of data from experiments with atmospheric torches provided the foundation for the general understanding of combustion grown diamond that is described in the following. A more comprehensive summary of experiments performed using atmospheric torches can be found in an earlier review [11].
B.
General Considerations
1.
Chemistry
The composition of acetylene flames is much different from the traditional H2 /CH4 mixtures used in hot-filament and plasma CVD systems. However, the chemistry and the mechanism of diamond growth in combustion CVD
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are expected to be similar to those of other diamond growth systems. First, at R = 1.0 the acetylene flame is positioned in the direct center of the empirical growth regime diagram developed by Bachmann et al. [12] (see Chapter 4). The detailed chemistry of acetylene flames is discussed in detail in Sec. IV, but its essence may be distilled to the following two reactions: C2H2 ⫹ O2 → 2CO ⫹ H2 H2 ↔ H ⫹ H
⌬H⬚R = ⫺107 kcal/mol
(1)
Reversible: function (T, P)
(2)
First, nearly all of the oxygen and acetylene is rapidly consumed, producing carbon monoxide and hydrogen as described by the global reaction (1) [13]. The highly exothermic nature of this reaction results in flame temperatures >3000 K. At these temperatures a significant fraction of hydrogen molecules are converted into atomic hydrogen as expressed by global reaction (2) [14]. The degree of dissociation is strongly dependent on both temperature and pressure. Carbon monoxide is thermally stable and not expected to affect the growth processes under these conditions. A small fraction of the acetylene is converted into hydrocarbon radical species or remains unreacted. Thus, growth schemes discussed in Chapter 4 that suggest that diamond growth proceeds through the addition of hydrocarbon radicals in the presence of a superequilibrium concentration of atomic hydrogen may also be applied to combustion systems. Where filament and plasma systems employ electrical energy to create atomic hydrogen, combustion systems rely on the chemical energy stored in the acetylene molecule to generate heat and atomic hydrogen. 2.
Substrate Temperature
Substrate temperatures during the combustion CVD of diamond range from 950 to 1650 K. Because of these high temperatures, substrates have been limited to materials such as molybdenum, alumina, silicon, and diamond. Evaluating the true substrate temperature is difficult in all diamond CVD environments but especially so in combustion synthesis due to the extreme heat fluxes present. The two techniques most commonly used are optical pyrometry and thermocouples placed on the back side of the substrate. Pyrometry provides a good measure of the relative temperature of the substrate, but absolute values may be off by as much as 100 K due to uncertainties in emissivities. Care must be taken to position the pyrometer in order to avoid emission from the flame itself. Back-side thermocouples are affected by both the quality of the contact and the steep thermal gradients present. Both techniques work well for a single configuration, but caution is urged when comparing absolute temperatures obtained from different systems.
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Substrate temperature has a dramatic effect on two important properties: growth rate and morphology. Several authors [15–17] have observed that the growth rate increases exponentially with temperature as shown schematically in Fig. 3. The behavior shown by the curve in Fig. 3 is characteristic of combustion systems; however, its location may be shifted to the right or left depending on the operating conditions and the technique used to measure temperature. Weimer et al. [15] measured activation energies between 12 and 18 kcal/mol for diamond growth on individual crystal faces. Slightly higher values (16–23 kcal/mol) have been reported for polycrystalline films [16,17]. These values are in good accord with activation energies measured in hot-filament systems, indicating that the growth mechanisms may be similar [18,19]. As the substrate temperature increases, the growth rate reaches a maximum (Fig. 3). Increasing the substrate temperature beyond this point causes a rapid decline in both the quality and the growth rate [15–17]. In atmospheric torches the maximum growth rate occurs at substrate temperatures between 1450 and 1650 K [15,20]. In lowpressure (40–50 torr) flat flame burners the same growth rate behavior has been observed, but the maximum occurs at significantly reduced temperatures of 1050–1250 K [17,21]. Kim and Cappelli [21] suggested that the falloff at high temperature may be due to oxidation of diamond. An alter-
Figure 3 The effect of substrate temperature on growth rate observed in combustion CVD of diamond. (From Refs. 15–17, 20.)
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native possibility is that the growth precursor desorbs from the surface before having the opportunity to incorporate into the diamond lattice [22]. Substrate temperature also determines the diamond film morphology. The morphology of most films deposited by combustion CVD are typically dominated by {111} facets [20,21,23]. However, a number of researchers [17,20,23,24] have observed that at higher temperatures the morphology consists of {100} facets parallel to the substrate. As temperature increases, the growth habit changes from octahedron to cubic as shown in Fig. 4. The temperature where {100} faceting occurs was near the temperature where maximum growth rates were observed (Fig. 3) [17,20]. 3.
Reactant Composition
In atmospheric torches it has been observed that morphology also depends on reactant composition [25,26]. Decreasing the oxygen fraction has an effect similar to that of increasing substrate temperature as shown in Fig. 4. It was observed that the morphology changed from octahedron growth under oxygen-rich conditions to cubic growth as the acetylene content was increased [25,26]. The dependence on composition is in accord with observations of morphology variations with methane concentration in plasma reactors [27,28]. For flat flame burners this dependence on composition has not been reported, perhaps because the range of R where diamond is successfully formed is much smaller [17,21]. Of the two variables, morphology depends more strongly on substrate temperature than composition. C.
Modifications of the Welding Torch
In the welding torch operation shown in Fig. 1, diamond is often deposited in a small circle or in an annular ring. To increase the deposition area a
Figure 4 Schematic diagram showing changes in crystal growth habit as a function of both substrate temperature and reactant composition. (From Refs. 17, 20, 23–26.)
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number of modifications of the welding torch have been investigated, including turbulent flame operation, multinozzle burners, and enclosed flames. A few groups [29,30] have investigated operating the welding torch under turbulent flow conditions. The turbulent condition was achieved by increasing both the flow rate and the burner orifice. Despite some reports [31] of improved quality, it was generally concluded that the changes were not dramatic and not worth the costs associated with increased flow rates [29,30,32]. Some success was achieved with a multinozzle configuration was developed by Zhu and coworkers [33], although the morphology was not uniform over the deposition area. During atmospheric operation, diffusion of oxygen or impurities from the ambient may effect film quality. Attempts to minimize the effect by surrounding the torch with a coannular flow of inert gas have been studied [34,35]. Under these conditions the outer flame (Fig. 1) was suppressed, but no dramatic changes were observed in quality or growth rate. Another way to exclude the ambient is simply to enclose the torch in a vacuum chamber [36–39]. It was observed that the outer diffusion flame disappeared as there was no oxygen present and the deposition area was increased by 20% [36,37]. In addition, reducing the pressure to 300 torr increased the deposition area by 200%, but the growth rate decreased significantly. Morrison et al. [38] performed a systematic study of an enclosed flame at 700 torr and found that a low-temperature pretreatment step enhanced the nucleation density. Wang et al. [39] achieved some success using rastering techniques with an enclosed flame to coat larger areas.
III.
PRACTICAL IMPLEMENTATION: THE FLAT FLAME BURNER
A.
The Flat Flame Burner
1.
Burner Design
A number of groups [14,17,21,40,41–45] have addressed the issues of increasing the deposition area and improving uniformity through the use of flat flame reactors operated at both atmospheric and reduced pressure. Cooper and Yarbrough [40] first demonstrated diamond deposition in a commercially available burner at pressures between 25 and 40 torr. A schematic of a flat flame burner CVD for reduced pressure operation is shown in Fig. 5. Premixed gas enters a water-cooled burner and exits through a matrix that creates a radial uniform velocity profile with only an axial component. In practice this has been achieved by drilling 1-mm holes in a copper plate [21,41], and by using a honeycomb made of thin-walled silica tubes
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Figure 5 pressures.
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Schematic of a flat flame burner CVD system operated at reduced
[17]. A flat, circular flame is formed stabilized below the burner surface as shown in Fig. 5. This is a very different structure than the conical flame that is associated with atmospheric torches (Fig. 1). The substrate is maintained at a certain temperature T and positioned at a distance L from the burner surface. This design was developed in the 1950s to measure flame velocities near extinction/ignition limits [46,47]. The burner’s characteristics of stability and geometrical simplicity that made it useful for that purpose are equally valuable for chemical vapor deposition. Atmospheric pressure flat flames are less stable and require more complicated burner design. Two examples that have been used include a coflow arrangement and a trumpet-bell design that are shown schematically in Fig. 6 [42]. Murayama and coworkers [43,44] used the coflow design with hydrogen as the external gas to stabilize their flame. Workers at Sandia have demonstrated the scalability of the trumpet-bell design [42,45]. Comparison of Figs. 5 and 6 shows another important difference between atmospheric and low-pressure operation. At low pressures the flame is stabilized on the burner surface, whereas atmospheric flames are substratestabilized flames located within a millimeter or two of the substrate. In lowpressure operation the substrate may be positioned at any distance from the burner, whereas in atmospheric systems the distance is constrained by the stability limits. In addition, in large-area atmospheric flames it was found that it was necessary to surround the substrate with equally spaced pillars to stabilize the flame [48].
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Figure 6 Two designs of atmospheric flat flame burners: (a) a coflow design and (b) a trumpet bell design. (From Ref. 48.)
2.
Radial Uniformity and Scalability
Regardless of operating pressure, the unique feature of this design is that films are produced that are uniform in both thickness and morphology. Figure 7 shows micrographs of surface morphology and a cross-sectional profile that demonstrate the radial uniformity achieved in flat flame burners. At low pressures, Glumac and Goodwin [41] demonstrated the scaleable nature of the system, producing uniform films over areas as large as 13 cm2. More recently, Hahn et al. [47] scaled up an atmospheric deposition system to 20 cm2. Another feature of the flat flame CVD reactor is that it is amenable to detailed kinetic modeling. Under proper operating conditions [49], the axisymmetric stagnation flow may be treated in one dimension using a similarity solution [50]. A number of groups have performed detailed kinetic modeling on these systems, which has provided insight into the deposition mechanisms. These studies are discussed in more detail in Sec. IV.
B.
Operating Conditions
1.
Reactant Composition
Low-pressure flat flame burners are operated with a value of R very near unity [17,21]. At atmospheric pressure, pure oxygen-acetylene flames are unstable. It has been found that the addition of hydrogen (18–25% by volume) while maintaining a C/O ratio ⬃1 was required to stabilize the flame and achieve diamond growth [42–45]. Because of the homogeneous nature
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Figure 7 Scanning electron micrographs (a–c) and cross-section profile (d) showing the radial uniformity of both morphology and growth rate achieved in flat flame burners. (From Ref. 17.)
of these flames, the composition range where diamond is deposited is much smaller than observed with the welding torch design [4]. Most researchers [42–45] report only a single composition that successfully produced highquality diamond. In low-pressure systems there is a small range where diamond is deposited [21,51]. The reason for the narrow range is shown in Fig. 8, where the equilibrium mole fractions of selected species are plotted as a function of composition. The three major species (CO, H2 , and H) are nearly constant across this range of R. For the minor species a very sharp transition is observed between R = 0.98 and R = 1.02. The concentration of carbon-containing species (C2H2 , CH3 , C) falls between four and seven orders of magnitude over this narrow regime. A concomitant increase in oxidation products (H2O, CO2 , O2) is also observed. Certainly, very good control of mass flow controllers is required for reproducible results using combustion synthesis. 2.
Operating Pressure
As already discussed, operating pressure has a major influence on burner design and flame stability. The pressure also influences a number of critical
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Figure 8 Equilibrium mole fractions of selected species in an oxyacetylene flame as a function of reactant composition R = O2 :C2H2 . Calculation performed at 40 torr with 5% argon.
issues including growth rate, temperature control, and deposition temperature. First, growth rates in atmospheric systems range from 25 to 40 m/hr [44,45]. In comparison, growth rates in flat flame burners at ⬃45 torr have been reported between 4 and 5.5 m/hr [17,21]. Although the gas density is about 15 times greater in atmospheric reactors, the growth rate is only a factor of ⬃7 times greater. More important with regard to economics, the carbon conversion efficiency (the fraction of acetylene atoms converted to diamond) is very similar in both cases. As discussed in Sec. II, growth temperatures in atmospheric systems are hundreds of degrees hotter than at low pressure. Substrate material selection is more limited for atmospheric flames, and these films are subject to the formation of large compressive stresses during cooling due to differing coefficients of thermal expansion [49]. Robust substrate temperature control is requisite for both controlling growth rate and achieving textured growth [17]. The proximity of the flame and substrate in atmospheric reactors re-
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Table 1 Comparison of Operating Flat Flames at Atmospheric and Reduced (40–50 torr) Pressures Aspect
Reduced pressure
Atmospheric pressure
Growth rates Carbon conversion efficiency Deposition temperatures Temperature control Flame stability Safety Vacuum chamber required Burner design
4–5.5 m/hr Similar ⬃0.005% 400–1000⬚C Easy Yes Advantage Yes Simple
25–40 m/hr Similar ⬃0.005% 700–1200⬚C Difficult No: requires H2 addition No Complex
Source: Refs. 17, 21, 42–45, 51, 52.
quires that energy must be removed at fluxes on the order of 500 W/cm2, requiring the use of elaborate cooling strategies [20,45,48]. In contrast, lowpressure reactors are nearly thermally neutral, and films have been deposited without any cooling present [41]. Simple control strategies can maintain the substrate temperature within a few degrees of a desired setpoint for hours of deposition [52]. Although atmospheric burner design is more complicated, an advantage is that no vacuum chamber is required. Table 1 summarizes some of the pros and cons of operating pressure in diamond deposition. C.
Applications
Combustion CVD was invented 7 years after the discovery of hot-filament and plasma techniques, and subsequently the majority of applied research has been performed in these conventional reactors. However, the studies that have been conducted demonstrate that combustion-grown diamond is comparable to material produced in other systems. Some of these examples are presented in this section with the intention of demonstrating the flexible nature of combustion CVD. 1.
Textured Film Growth
For a number of applications [53,54] films that demonstrate {100} texturing have several advantages. The combination of {100} growth conditions [27] and bias-enhanced nucleation techniques [55] has yielded highly oriented diamond films. Both the number and angle of grain boundaries are significantly reduced in these films relative to randomly oriented films [28,52]. Both the thermal [56] and the electronic [53] properties of diamond are
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adversely effected by the presence of grain boundaries. For example, the measured carrier mobility was three times greater in {100} oriented films than polycrystalline diamond [53]. In addition, {100} growth provides a means of planarizing the growth surface [52], enabling the use of conventional silicon technology to manufacture device structures without expensive grinding and polishing steps [54]. Figure 9 shows a scanning electron micrograph of a highly oriented {100} film that was produced by a two-step process. First, oriented nucleation was achieved using bias-enhanced nucleation techniques in a conventional microwave plasma reactor. The substrate was transferred to a low-pressure combustion reactor and grown for 3 hr [52]. In Fig. 9 one can see that a number of grains have coalesced, forming a very flat surface. 2.
Doping and Device Fabrication
Ravi and coworkers [57] demonstrated the high purity of combustion-grown diamond, producing films that had resistivities two orders of magnitude greater than natural diamond. The doping of diamond imparts semiconductor properties that are required for a number of applications [53]. Most of the limited work on doping of combustion-grown diamond has been carried out at the Naval Research Laboratory [58–61]. Glesener and coworkers [58]
Figure 9 Scanning electron micrograph of a highly oriented {100} film formed by a combination of bias-enhanced nucleation in a microwave reactor and textured growth in a low-pressure combustion CVD system. (From Ref. 52.)
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fabricated Schottky diodes by depositing boron-doped diamond on molybdenum. Further studies [59,60] indicated that high surface resistance was critical to form good junctions. The boron was added by bubbling part of the acetylene through a solution of diboric trioxide in acetone. Doverspike and coworkers [61] used an atomizer to entrain aerosol droplets that contained dopant species into the gas flow. Boric acid dissolved in methanol was used as a dopant solution. Recently, boron incorporation has been achieved without the use of a solvent. Trimethylborate is a high-vapor-pressure liquid that may be introduced by mass flow controllers at low pressure or with a bubbler for atmospheric operation [14]. 3.
Low-Temperature Deposition
Low-temperature processing is required for compatibility with a number of optoelectronic materials. In addition, low-temperature deposition minimizes the compressive stress that is formed when diamond is grown on material with differing coefficients of thermal expansion. Diamond deposition at rates of ⬃1 m/hr on glass substrates was achieved at low temperature (750 K) in a low-pressure combustion reactor [51]. The low temperature was achieved by increasing the burner-substrate distance to 20 mm. Continuous, adherent films were deposited on Vycor and Pyrex glasses. The best films were optically transparent and no damage to the substrates was observed.
IV.
DEPOSITION CHEMISTRY
A.
Diagnostics
The chemistry of combustion diamond CVD has been primarily investigated using laser-induced fluorescence (LIF) and mass spectroscopy. The two techniques are complementary. Mass spectroscopy has been used to provide quantitative measurements of the major stable species, which are important for verifying kinetic models. In this technique a quartz microprobe samples gases at the diamond growth surface. Using mass spectroscopy, Wolden et al. [62] noted that the strong compositional dependence observed in a flat flame reactor was related to a sharp change in the acetylene concentration. On the other hand, LIF is a nonintrusive optical technique that has been used to detect radicals such as C2 , C2H, CH, and OH in the gas phase above the substrate. Researchers using plasma and hot-filament reactors have focused on C1 species such as methyl radical and atomic carbon as the dominant growth species [63,64]. However, evidence from combustion reactors suggests that C2 species are also very important. Three different groups [13,65,66] have noted a strong relation between C2 concentration profiles
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and diamond deposition rates. The role of CH is less clear, and OH appears to be of minor importance [65]. The work of Cappelli and Paul [65] suggested that C2H also plays an important role in the deposition process. B.
Reactor Modeling
An attractive feature of flat flame reactors is that they may be simulated using simple stagnation flow models with complex combustion kinetic mechanisms. Several workers have used this approach to model flat flames at diamond growth conditions [66–68]. Kim and Cappelli [67] argued that optimum diamond growth conditions result from a competition between hydrocarbon growth and oxidation. Simulations at Sandia indicated the importance of operating at high flame speeds to increase atomic hydrogen production [68]. A detailed kinetic mechanism containing 50 species and 218 reactions has been developed for the acetylene flame by Miller and Melius [69]. The model has been verified experimentally by Glumac and Goodwin [70] using mass spectroscopy and LIF at conditions slightly oxygen rich for diamond growth. Wolden and coworkers [14] performed a similar validation study at diamond growth conditions with a diamond substrate present. Figure 10 shows the profiles of selected species between the burner and substrate at diamond deposition conditions [14]. Acetylene and oxygen are rapidly converted to hydrogen and carbon monoxide as the gas exits the burner. At the high temperatures involved, a large amount of atomic hydrogen is produced. The concentration of H decreases near the surface due to recombination reactions with the diamond surface [71,72]. As noted in Sec. III, there is a very narrow composition window where diamond growth is achieved. Figure 11 shows the normalized concentration of selected species from model predictions plotted over a range of composition that spans the transition from graphite deposition to diamond deposition through etching [66]. The two species most often implicated in diamond growth, H and CH3 , hardly changed over this range. In contrast, the acetylene concentration decreased by 70%, which was the largest change observed for any species [66]. Although the mechanism is not fully understood, acetylene appears to be responsible for the sharp transition between diamond and graphite deposition.
C.
Alternative Fuels
Although acetylene is the most common fuel used for combustion synthesis, a number of alternatives have been examined including hydrogen [73], eth-
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Figure 10 Profiles of selected species between the burner (11 mm) and the substrate (0 mm) at diamond growth conditions: R = 1.00, P = 47 torr, Ts = 750⬚C, Uo = 600 cm/sec. (From Ref. 14.)
ylene [74], MAPP gas [75],* and methane [76]. The interest in other fuels is economic, due to the relatively high cost of acetylene. Diamond has been deposited successfully in each case, but the growth rates are significantly less than with acetylene. The superiority of acetylene is due to its high flame velocity and temperature. Table 2 compares these properties for a number of fuels. Examining Table 2, one finds that the flame temperature of acetylene is at least 350⬚C hotter than for the alternatives. All of these fuels produce hydrogen, but only atomic hydrogen is useful for diamond deposition. The importance of acetylene’s high flame temperature is shown in Fig. 12, which shows the dependence of hydrogen dissociation on both temperature and pressure. First, the temperature dependence is exponential. Thus, increasing the temperature from 2500 to 2900 K results in a fourfold increase in the H atom mole fraction. Besides growth rate considerations, modeling efforts
*MAPP gas is a commercially available mixture of about half liquefied petroleum gas (mostly propylene) and half C3H4 (a mixture of the two isomers methyl acetylene and propadiene).
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Figure 11 Variations of selected species as a function of reactant composition R = O2 :C2H2 over a range that spans the transition from graphite deposition to diamond-deposition to etching. (From Ref. 66.)
Table 2 Comparison of Adiabatic Flame Temperatures and Flame Speeds of Selected Fuels Fuel Acetylene, C2H2 Ethylene, C2H4 Hydrogen, H2 Propylene, C3H6 Methane, CH4
Adiabatic flame temperature (K)*
Flame speed (cm/sec)
2910 2565 2525 2505 2325
141.0 68.3 264.8 43.8 33.8
*Assumes base temperature = 25⬚C; oxidizer, air; equivalence ratio, 1.0.
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Figure 12 Equilibrium dissociation of atomic hydrogen as a function of temperature and pressure.
suggest that diamond quality increases quadratically with atomic hydrogen concentration [77]. Figure 12 also indicates that H-atom dissociation is enhanced at lower pressures in accordance with Le Chaˆtelier’s principle.
V.
COMPARISONS AND CONCLUSIONS
Although material deposited by combustion CVD has not been tested extensively, there is every indication that the quality is on par with that of diamond produced in hot-filament and plasma systems. Schottky diodes and high-energy particle detectors have been successfully fabricated using combustion-grown diamond. Issues of large-area deposition and uniformity have been addressed with flat flame burners. Growth rates are exceeded only by atmospheric plasma torches. Important features such as textured growth, boron doping, and low-temperature deposition have all been achieved in combustion reactors. Oriented nucleation, which has been achieved by substrate biasing in microwave and hot-filament reactors, has yet to be demonstrated in a combustion CVD system.
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Early reports concluded that combustion CVD was a much more expensive synthesis technique than plasma or hot-filament methods [78]. However, more recent analyses using data from a low-pressure flat flame burner indicates that the costs are very similar to those of microwave plasma CVD (M. Stonefield and C.A. Wolden, unpublished work). In comparison with plasma-aided processed, combustion CVD requires much less capital but has higher consumable costs because of the amounts of acetylene and oxygen required. In conclusion, the status of combustion synthesis of diamond using acetylene flames has been reviewed. The development of the technique from its inception using conventional welding torches to its present implementation using flat flame burners has been detailed. High growth rates and uniformity are characteristic of diamond deposition in flat flame burners. The design may be scaled to arbitrarily large areas. The quality of combustion diamond is comparable to that of material produced by other techniques. Economically, combustion CVD is competitive with plasma techniques. Combustion CVD requires less capital investment, but these savings are offset by high material costs. Substantial savings will be realized by improvements in growth rates and carbon utilization efficiency.
ACKNOWLEDGMENTS The author acknowledges the financial support of the National Research Council and Army Research Office during the writing of the manuscript. The author would like to thank Dr. John Prater of the Army Research Office for many helpful discussions and for stimulating my work in combustion CVD of diamond. Collaborations with the research groups of Professors Zlatko Sitar and Robert Davis in the Materials Science and Engineering Department at North Carolina State University were greatly appreciated.
REFERENCES 1. 2. 3. 4. 5. 6. 7.
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Wolden XH Wang, W Zhu, J von Windheim, JT Glass. J Cryst Growth 129:45, 1993. DB Oakes, JE Butler, KA Snail, WA Carrington, LM Hanssen. J Appl Phys 69:2602, 1991. LM Hanssen, KA Snail, WA Carrington, JE Butler, S Kellog, DB Oakes. Thin Solid Films 196:271, 1991. PW Morrsion Jr, JT Glass. In: G Davies, ed. Properties and Growth of Diamond. London: INSPEC, 1994, p 380. PK Bachmann, D Leers, H Lydtin. Diamond Relat Mater 1:1, 1991. Y Matsui, A Yukki, M Sahara, Y Hirose. Jpn J Appl Phys 28:1718, 1989. CA Wolden, RF Davis, Z Sitar, JT Prater. Diamond Relat Mater 7:133, 1998. RA Weimer, TP Thorpe, KA Snail. J Appl Phys 77:641, 1995. KA Snail, CM Marks. Appl Phys Lett 60:3135, 1992. CA Wolden, Z Sitar, RF Davis, JT Prater. Appl Phys Lett 69:2258, 1996. E Kondoh, T Ohta, T Mitomo, K Ohtsuka. Appl Phys Lett 59:488, 1991. C Chu, M D’Evelyn, R Hauge, J Margrave. J Appl Phys 70:1695, 1991. JJ Schermer, JEM Hogenkamp, GCJ Otter, G Janssen, WJP van Enckevort, LJ Giling. Diamond Relat Mater 2:1149, 1993. JS Kim, MA Cappelli. J Mater Res 10:149, 1995. CA Wolden, KK Gleason. Diamond Relat Mater 5:1503, 1996. GH Ma, Y Hirose, S Amanuma, M McClure, JT Prater, JT Glass. In: R Messier, JT Glass, JE Butler, R Roy, eds. New Diamond Science and Technology. Pittsburgh: Materials Research Society, 1991, p 587. KV Ravi, A Joshi. Appl Phys Lett 58:246, 1991. KV Ravi, CA Koch, HS Hu, A Joshi. J Mater Res 5:2356, 1990. K Hirabayashi, Y Hirose. Diamond Relat Mater 5:48, 1996. C Wild, R Kohl, N Herres, W Muller-Sebert, P Koidl. Diamond Relat Mat 3: 373, 1994. BR Stoner, SR Sahida, JP Bade, P Southworth, PJ Ellis. J Mater Res 8:1334, 1993. KA Snail, RG Vardiman, JP Estrera, JW Glesener, C Merzbacher, CJ Craigie, CM Marks, R Glosser, JA Freitas Jr. J Appl Phys 74:7561, 1993. JJ Schermer, LJ Giling, P Alers. J Appl Phys 78:2376, 1995. KA Snail, CL Vold, CM Marks, JA Freitas Jr. Diamond and Relat Mater 1: 180, 1992. JJ Scheimer, WA Elst, LJ Giling. Diamond and Relat Mater 4:1113, 1995. W Zhu, BH Tan, J Ahn, HS Tan. J Mater Res 30:2130, 1995. RC Aldredge, DG Goodwin. J Mater Res 9:80, 1994. T Abe, M Suemitsu, N Miyamoto. J Cryst Growth 143:206, 1994. M Murakawa, S Takeuchi. Surf Coat Technol 54/55:403, 1992. K Komaki, M Yanagisawa, I Yamamoto, Y Hirose. Jpn J Appl Phys 32:1814, 1993. PW Morrsion, A Somashekhar, JT Glass, JT Prater. J Appl Phys 78:4144, 1995. DY Wang, YH Song, JJ Wang, RY Cheng. Diamond Relat Mater 2:304, 1993. JA Cooper Jr, WA Yarbrough. Diamond Opt SPIE 1325:41, 1990. NG Glumac, DG Goodwin. Mater Lett 18:119, 1993. KF McCarty, E Meeks, RJ Kee, AE Lutz. Appl Phys Lett 63:1498, 1993.
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9 Laser-Assisted and Optical Pumping Techniques for Diamond Synthesis Vish V. Subramaniam and Shashi M. Aithal The Ohio State University, Columbus, Ohio
I.
INTRODUCTION
Diamond growth by chemical vapor deposition (CVD) processes requires activation of the gas phase adjacent to a substrate. The previous chapters of this book have discussed several methods of providing this activation needed for driving chemical reactions in the gas phase. Among these are hot-filament CVD (HFCVD), flame synthesis, and plasma CVD. All three methods rely upon thermal heating or near-equilibrium effects to induce reactions in the gas leading to diamond growth. Consequently, it is not possible to exercise selectivity from among the various competing chemical processes for diamond growth. Selectivity of desired chemical channels or lowering of growth temperatures, for instance, requires that attention be devoted to how the growth medium is energized and reacted. Lasers offer a unique means of activating the molecules in the growth medium. They are potentially capable of reducing growth temperatures as well as providing selectivity in a reactive process. In this chapter, the use of lasers in diamond synthesis is discussed. Use of lasers for diamond synthesis has not always been motivated with the goal of commercialization. Rather, lasers have been used to create environments that enable detailed study of the gas (or condensed) phase chemistry. These can also lead to better understanding of some of the reacting channels occurring in the other, more commercially viable processes 325
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leading to process improvements. There are also other reasons that laserbased methods of diamond synthesis have been explored. These include the need to reduce substrate and gas (or liquid) temperatures during growth and to control and limit the codeposition of nondiamond carbon by selectively driving only the specific chemical channels that lead to diamond formation. In addition, lasers can be used to control diamond deposition patterns and enhance growth rates. Processes that meet these requirements are not necessarily the most economical as well. Hence, cost is not the only issue to consider when comparing laser-based methods with other better established CVD processes and General Electric’s high-pressure, high-temperature (HPHT) process. In order to appreciate the utility of lasers in CVD processes for diamond growth, it is necessary to understand how the energy added to a gas is redistributed within the various modes of its molecules. Gases comprise molecules and atoms and, in the case of plasmas, ions and electrons as well. These molecules, atoms, and ions in turn possess various internal modes of energy storage in addition to their translational (or what is called external) modes of motion. Figure 1 shows these various modes schematically. In addition to the three external degrees of freedom (i.e., translation in the three coordinate directions) that a particle possesses, molecules and molecular ions can store energy in rotation, vibration, and electronic excitation. Atoms and atomic ions can store energy only in electronic excitation in addition to their external modes. When energy is added to a gas in the
Figure 1 Schematic showing the different modes of energy storage for a diatomic molecule.
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327
form of simple heating, it first appears in the translational motion of the gas particles. Temperature is the common measure used to describe the average translational energy of the gas. The energy added by heating eventually makes its way into the other modes of molecular motion via collisions. Usually, exchange of translational energy between molecules takes place within a few collisions, and equilibration of energy between rotational and translational modes usually requires on the order of tens of collisions. In contrast, many collisions (on the order of thousands or many more) are required for the vibrational modes of molecules to equilibrate with rotational and translational modes [1–3]. Because the excitation of the vibrational modes is helpful in driving endothermic reactions, it can be seen that simple thermal heating is a circuitous and inefficient route to cause a gas to react. Thermal stimulation of a gas can be accomplished in a number of ways, including simple heating, electrical heating by partially ionizing it and then driving a current through the gas, or using a laser. It is usually inefficient and expensive to use a laser for thermal heating of a gas unless repetitive, very short duration pulse heating is required for a particular process. Lasers are also useful in instances where only the gas adjacent to a surface must be heated without heating the surface itself (see Fig. 2). In cases where the surface alone needs to be heated to drive a heterogeneous reaction, the gas must be transparent to the laser source whose energy is absorbed at the substrate-gas interface. This is often useful in the case where spot heating as opposed to bulk heating is required. Lasers drive chemical reactions by exciting specific transitions in a molecule. It may be pertinent to ask at this point whether or not a laser is necessary to excite a specific transition in a molecule. In other words, can one use an intense lamp (with an appropriate filter) to produce monochromatic (but incoherent) light to initiate reactions in gases? The answer of course is yes, and chemists were employing such techniques long before the invention of lasers to measure reaction rates [4]. However, inference of re-
Figure 2 Schematic showing the use of a laser to drive homogeneous gas-phase reactions by pyrolysis, photolysis, or laser excitation without heating of the surface.
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action rates in complex chemically reacting systems is rendered difficult for two important reasons. First, incoherent monochromatic sources such as intense lamps are limited in intensity because they are diffuse, i.e., emit light in all directions. Usually, the higher the intensity (i.e., energy per photon multiplied by the number of photons per unit area, per unit time), the higher is the reaction yield, which makes it easier to measure changes in concentrations. Second, production of extremely short duration pulses (say of duration nanoseconds to hundreds of nanoseconds) of sufficient energy to drive fast reactions is nearly impossible with an incoherent, monochromatic source. Lasers, on the other hand, are monochromatic sources that can produce high pulse energies with extremely short durations and sustain these characteristics over long distances because they are coherent. We now digress a bit and examine this property of coherence in some detail, following the development in Ref. 5. Two sources are said to be coherent if they emit light or photons of a given frequency (or energy h, where h is Planck’s constant), with nearly the same amplitude, and in phase with each other. An important consequence of coherence is that a maximum intensity is observed at a point in space where the phase difference between two waves or photons reaching that point is an integral multiple of 2. Another way of stating this idea is that the maximum intensity will occur when the path difference1 between two waves or photons is an integral multiple of the wavelength (= c/, where c is the speed of light). The minimum intensity occurs when the phase difference between the photons is an odd multiple of , i.e., (2m ⫹ 1), where m is an integer (= 0, 1, 2, . . .). Coherent sources thus exhibit the important property of interference. The appearance of fringes (i.e., alternating segments of maximum intensity and minimum intensity) is due to interference. Because light is a form of electromagnetic radiation, it can be repre→ sented at any point in space as a time-varying electric field E(t) = → → E0e ⫺ite i(t), where E0 is the amplitude, = 2 is the circular frequency, i = 兹⫺1, t is time, and →(t) is the time-dependent phase. Thus, when two → photons represented by E1 and E2 interact, the resulting intensity can be expressed as a time average: →
→
→
→
→
→
I = 具E⭈E*典 = 具(E1 ⫹ E2)⭈(E* 1 ⫹ E*)典 2 →
→
→
→
= 具兩E1兩2 ⫹ 兩E2兩2 ⫹ 2 Re(E1 ⭈E*)典 2 where the angle brackets denote the time average, defined as 1
If the path difference is x and the phase difference is ␦, the two are related by ␦ = 2x/, where is the wavelength.
Laser-Assisted and Optical Pumping Techniques
具 f 典 = lim → t0
⬁
1 t0
冕
329
t0
f(t) dt
0
The symbol Re denotes the real part of a complex number, and the superscript * denotes the complex conjugate of a complex number. Now for simplicity, if we assume that the two fields are polarized,2 then their vectorial nature can be ignored. The intensity then becomes I = I1 ⫹ I2 ⫹ 2 Re具E1E*典 2 where I1 = 具兩E1兩2典 and I2 = 具兩E2兩2典. Thus, if t is the time required for one photon to travel via a path A and if (t ⫹ ) is the time required for a second photon to travel via path B, then the intensity due to interference is ⫹ )典 = 2 Re(⌫12()). The function ⌫12(t) is called the mutual 2 Re具E1(t)E*(t 2 coherence function or correlation function of the two fields. Analogously, ⌫11(t) is called the autocorrelation function. By definition, ⌫11(t = 0) = I1 and ⌫22(t = 0) = I2. The intensity is also written as I = I1 ⫹ I2 ⫹ 2兹I1I2 Re{␥12()} where ␥12() is called the degree of partial coherence and is given by
␥12() =
⌫12() ⌫12() = ⌫ (0)⌫ (0) 兹 11 兹I1I2 22
If 兩␥12兩 = 1, we have complete coherence. If 兩␥12兩 = 0, there is complete incoherence, and if 0 < 兩␥12兩 < 1, there is partial coherence. Now, suppose that the interference pattern varies between the two limits set by Imax and Imin: Imax = I1 ⫹ I2 ⫹ 2兹I1I2兩␥12兩 Imin = I1 ⫹ I2 ⫺ 2兹I1I2兩␥12兩 The fringe visibility is then defined as ᭙=
Imax ⫺ Imin 2兹I1I2兩␥12兩 = Imax ⫹ Imin I1 ⫹ I2
When I1 = I2, then ᭙ = 兩␥12兩. Thus, for complete coherence, ᭙ = 1 and the fringes have a maximum contrast of unity, whereas for complete incoherence, ᭙ = 0 and the fringe contrast is 0 (i.e., there are no distinguishable fringes). We will now attempt to relate the degree of partial coherence to the characteristics of a source. Suppose the source emits an electromagnetic 2
Simply defined, polarization is a chosen direction of the electric field vector in the plane normal to the direction of propagation of the electromagnetic wave.
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wave whose field varies sinusoidally for a certain duration 0 and then changes phase abruptly. This can represent, for example, the case of an excited atom or molecule that emits radiation for a certain amount of time and then suffers a collision with another atom or molecule. This sequence of emission for a certain duration 0 followed by an abrupt change of phase is then repeated indefinitely. 0 is called the coherence time. The phase therefore may vary in a manner such as that shown in Fig. 3 for illustrative purposes. Suppose now that the light from this source represented by E(t) = E0e⫺ite i(t) is split into two beams and subsequently brought together to produce interference. Assuming that 兩E1兩 = 兩E2兩 = 兩E兩, the degree of self partial coherence is
␥() =
具E(t)E*(t ⫹ )典 具兩E兩2典
Substituting for E(t), we have
␥() = 具e ie i [ (t)⫺ (t⫹)]典 = ei lim → T
⬁
1 T
冕
T
e i [ (t)⫺ (t⫹)] dt
0
Figure 4 shows an illustrative plot of the phase difference (t) ⫺ (t ⫹ ) versus time. For the first time interval, 0 < t < 0, (t) ⫺ (t ⫹ ) = 0 when 0 < t < 0 ⫺ . However, for 0 ⫺ < t < 0, the phase difference takes on some random value ⌬ between 0 and 2. Therefore, the degree of self partial coherence becomes
␥() = e
i
= e i
1 0
冕
0
e
i [ (t)⫺ (t⫹)]
dt = e
0
i
1 0
再冕
0⫺
0
冕
0
dt ⫹
0⫺
冎
e i⌬ dt
1 {0 ⫺ ⫹ e i⌬} 0
Each interval will have the same expression for ␥(), except that the value
Figure 3 Plot of phase (t) versus time t for the example of a quasi-monochromatic source. (From Ref. 5.)
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Figure 4 Plot of phase difference, (t) ⫺ (t ⫹ ) versus time t for the discussion of coherent length and coherent time. (From Ref. 5.)
of ⌬ will be different. Because ⌬ is random, the time average of e (i⌬) is zero, and the average of ␥() over all such intervals will be (1 ⫺ (/0)e i, whenever < 0. On the other hand, when > 0, the phase difference is always random so that its average over many intervals is zero. Hence, we have
␥() =
or
再 再
兩␥()兩 =
冉 冊
for < 0
0
for ⱖ 0
1⫺
1⫺ 0
e i 0
0
for < 0 for ⱖ 0
This function is shown schematically in Fig. 5. Since 兩␥兩 is equal to the fringe visibility ᭙ for two light waves of equal amplitude, it can be seen from Fig. 5 that ᭙ drops to zero when exceeds the coherence time 0. This means that the path difference between the two split light waves, say x, must not exceed the value of Lc = c0, where c is the speed of light. The quantity Lc = c0 is called the coherence length. Thus, for x > Lc, interference fringes will not be observed. Interference fringes can be observed only for x < Lc. In a radiating gas, the time between collisions of gas molecules is not constant but varies randomly from one collision to the next. However, an average time between collisions can be defined. Hence, 0 should be interpreted as an average value of the individual coherence times. The same holds true for the coherence length.
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Figure 5 Plot of degree of self partial coherence, 兩␥ ()兩 versus for the discussion of coherent length and coherent time. (From Ref. 5.)
We are now in a position to understand the crucial differences between lasers and ordinary, intense, incoherent sources. The first point to realize is that no source of light is ever strictly monochromatic. There is always a spread of frequency about some mean frequency due to natural broadening.3 It can be shown using Fourier analysis that the width ⌬ of the frequency distribution is given by ⌬ = 2/0 or ⌬ = 1/0 [5]. Conversely, we can also say that a spectral line of width ⌬ has a corresponding coherence time of 0 = 1/⌬, with a coherence length of Lc = c0 = c/⌬. This condition can also be expressed in terms of wavelength as Lc = 2/⌬ because ⌬/ = 兩⌬兩 /. For ordinary light sources such as lamps or discharge tubes, the line widths in the visible portion of the spectrum (say around 500 nm) are on the order of an angstrom [5]. The corresponding coherence length is then on the order of 5000 wavelengths or about 2.5 mm. This means that with such light sources, interference fringes would disappear beyond a distance of about 2.5 mm. In contrast, the line width of a gas laser can be as narrow as 103 Hz or less,4 which corresponds to a coherence length of about 300 km! Thus, coherent sources are capable of producing interference effects over much larger distances than incoherent sources. The other essential difference between incoherent and coherent light sources is the property of directionality. Ordinary lamps and discharge tubes, which produce incoherent light, radiate in all directions. In contrast, light
3
Natural broadening refers to the finite line width (centered around a specific wavelength) observed for a spectral line due to the finite radiative lifetime of a given excited state of an atom or molecule. The width associated with natural broadening has its origins in Heisenberg’s uncertainty principle, which gives a relationship between uncertainties in time and energy. 4 The corresponding line width for ordinary, incoherent light sources is on the order of 1010 Hz.
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from a laser can come very close to an ideal plane wave whose divergence is primarily due to diffraction effects [6]. Therefore, the high directionality of the laser beam can be used to focus it to extremely small dimensions and thus attain high power densities or intensities. For the aforementioned reasons, laser-based methods have a great potential to drive chemical reactions. It is more efficient to stimulate internal modes of gas molecules (i.e., the more reactive modes such as vibrational or electronic excitation) directly rather than relying upon thermal heating. The former relies upon sustaining thermodynamic disequilibrium between internal and external modes of molecular motion. Excitation of the internal modes with minimal excitation of the translational modes also leads to some control over selectivity, as we shall see later in this chapter. In contrast, thermal heating of a gas, results in states that are in thermodynamic equilibrium or in local thermodynamic equilibrium (LTE).5 Lasers, depending on how they are used, are capable of producing gaseous environments that can be in LTE or far from it. Processes that use this departure from LTE to induce selective growth of diamond or produce films of desired properties or characteristics are discussed later in this chapter. Their ultimate utility in industrial applications, though, must be evaluated with regard to cost.6 The use of lasers for diamond synthesis (or synthesis of other materials for that matter) can be grouped into general categories: (1) rapid heating and pyrolysis, shown schematically in Figs. 2 and 6; (2) ablation or physical vapor deposition (PVD), shown schematically in Fig. 7; (3) photolysis, shown schematically in Figs. 2 and 6; (4) laser excitation or optical pumping, shown schematically in Figs. 2 and 6; and (5) hybrid methods. An obvious use of a laser is for intermittent, rapid heating. Laser pyrolysis is included in this grouping, although there are some important differences. A laser can be used to heat a gas by irradiating surfaces in contact with the gas. This takes advantage of the fact that solid surfaces tend to absorb radiant energy broad band (i.e., over a continuous range of wavelengths) whereas gases tend to absorb discrete lines (i.e., at specific wavelengths). Such an instance of a solid surface being heated is encountered 5
Strictly speaking, the state of thermodynamic equilibrium is defined as one that is homogeneous, isotropic, and devoid of any gradients. However, most practical engineering situations involve flows and gradients. Therefore, the concept of local thermodynamic equilibrium or LTE is typically introduced as an approximation to allow a flowing gas exposed to gradients to be treated as consisting of infinitesimal regions that are nearly homogeneous, isotropic, and devoid of gradients. 6 We shall address the issue of cost later, as each technique is described. We hasten to point out that the economics of diamond is largely artificial, and therefore we will avoid using units of dollars per carat to discuss cost. Rather, we evaluate processes based upon energy cost expressed here in units of joules per carat ( J/carat).
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Figure 6 Schematic showing the use of a laser to drive gas-phase or heterogeneous reactions by pyrolysis, photolysis, or laser excitation while heating the surface locally.
when heterogeneous (i.e., gas-solid) reactions need to be driven. The laser wavelength must then be chosen so that the gas is transparent to this wavelength while the surface is not. A potential advantage over other methods of heating the surface is that the laser energy can be absorbed at the surface alone, with minimal heating of the bulk material. If the gas alone is to be heated, the laser energy must first be absorbed7 by the gas (at a specific wavelength or discrete wavelengths corresponding to vibrational or electronic transitions in the gas molecules). The energy initially absorbed in the internal modes is then rapidly redistributed to the external (translational) modes, thus raising the gas temperature and driving chemical reactions. This process, known as pyrolysis, can involve either specific constituents or all constituents of a gas mixture. Lasers can also be used to irradiate and vaporize a target material, with the vapor-phase products subsequently condensing on a surface placed near the target. We shall refer to this technique as laser ablation, and a schematic is shown in Fig. 7. In the literature, laser ablation is also known as laser physical vapor deposition or LPVD. This method has been used successfully to produce thin films of superconductors and other materials. The target or targets are usually pure materials or materials of fixed, known composition. Intense laser energy (usually in the ultraviolet) is focused on the target,
7
Laser energy can be absorbed by a gas molecule M either by resonant single-photon absorption or by multiphoton absorption. The former utilizes transitions between real energy levels in a molecule whereas the latter can also occur via virtual energy states: single-photon: h ⫹ M(E1) → M(E2) where E2 > E1 multiphoton: nh ⫹ M(E *) → M(E *) where E * 1 2 2 > E* 1 and E * 1 and E * 2 do not have to be real energy states belonging to the molecule M
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Figure 7
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Schematic of laser physical vapor deposition or LPVD process.
causing portions of it to vaporize and in some cases ionize or react. The vapor then is made to condense on a substrate placed nearby (which may be heated or cooled), resulting in a thin film. In most cases, the laser also interacts with the vapor (which may or may not be ionized) emanating from the target. Thus, the laser ablation technique could be easily applied in conjunction with other techniques such as plasma CVD, hot-filament CVD, or flame synthesis, resulting in a hybrid method. The third category is photolysis, wherein multiphoton processes are used to break apart polyatomic molecules into smaller, radical fragments. Multiphoton absorption is a nonlinear process whereby two or more photons are absorbed by a molecule, resulting in dissociation of the polyatomic molecule into molecular or atomic fragments. The process can be written symbolically in the following form of a chemical reaction: Mn ⫹ p(h) → Ra ⫹ Sb
(1) ⫺34
where p photons of frequency are absorbed (h = 6.626 ⫻ 10 J-sec, is Planck’s constant) by a polyatomic molecule containing n (n = a ⫹ b) atoms. The result is that the molecule is taken to its dissociation limit and the resulting polyatomic fragments are R (containing a atoms) and S (containing b atoms). The two ways in which a multiple number of photons can be absorbed by a polyatomic molecule are shown schematically in Fig. 8. The photons do not necessarily have to access real energy levels within the molecule en route to the dissociation limit. The multiphoton process can access virtual energy states [4,7,8]. Further, the absorption process can be either sequential or simultaneous. An example of a naturally occurring multiphoton process is the dissociation of O2 in the earth’s upper atmosphere due to ultraviolet (UV) radiation from the sun: O2 ⫹ m(h) → O ⫹ O
(2)
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Figure 8 Schematic showing multiphoton dissociation (MPD) process in a polyatomic molecule. MPD processes can take place through interaction of the molecule with many photons either sequentially or simultaneously and can involve either real energy levels or virtual energy levels.
where m is at least 2 for photons with a wavelength of 193 nm.8 The O atoms subsequently recombine with O2 to form ozone (O3). This technique of photolysis is routinely employed by chemists to measure reaction rates involving specific molecules or radicals [4,7,8]. There are laser-based techniques that take advantage of the coincidence between the wavelength of the laser light and energy difference for a particular transition (either electronic or vibrational) in a given molecule. Because such an application typically does not involve high powers but rather the right wavelength, these techniques are collectively referred to here as laser excitation or laser-excited CVD. A listing of different lasers presently available along with their operating wavelengths is given for reference in Table 1. A broader term that encompasses incoherent light (such as UV lamps) as well as coherent (laser) light is optical pumping. The terms laser excitation and optical pumping will be used interchangeably throughout this chapter, even when a laser is used to create the optical pumping. In optical pumping of a gas, the absorption is usually by a resonant single-photon process:
8
The particular wavelength of 193 nm is also the wavelength at which the ArF (argon-fluoride) excimer laser lases. The wavelength of 193 nm is in the range of absorption for O2 via the Schumann-Runge bands [4].
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Table 1 Available Laser Sources and Their Respective (Commonly Found) Wavelengths Laser type
Wavelength
ArCl excimer ArF excimer KrCl excimer KrF excimer Frequency-quadrupled Nd:YAG XeBr excimer XeCl excimer Nitrogen laser XeF excimer 7-Hydroxycoumarin dye 4-Methylumbelliferone dye Esculin dye Argon ion laser Argon ion laser Frequency-doubled Nd:YAG Puronin B dye Na-fluorescein dye 2,7-Dichlorofluorescein dye Rhodamine 6G dye Acridine red dye Rhodamine B dye Helium-neon laser Ruby (Al2O3 doped with Cr2O3) Nd:YAG (Y3Al5O12) HF chemical laser DF chemical laser CO laser CO2 laser Free electron laser
175 nm 193 nm 222nm 248 nm 266 nm 282 nm 308 nm 330 nm 351 nm 450–470 nm 450–470 nm 450–470 nm 488 nm 514.5 nm 532 nm Yellow 530–560 nm 530–560 nm 570–610 nm 600–630 nm 605–635 nm 632.8 nm 692.9 nm and 694.3 nm 1.064 m (dominant line) 2.7 m 3.8 m Multiple lines from 4.7 to 5.8 m 9.4 m, 10.6 m Visible to ultramicrowave (mm) wavelengths
Mn(E1) ⫹ h → Mn(E2)
(3)
where h = E2 ⫺ E1, and E1 and E2 are energy levels within the molecule Mn, comprising n atoms. This is followed by collisional processes that lead to energy transfer and subsequent chemical reaction. Lasers have been used successfully to synthesize diamond particles, films, and diamond-like films. Although lasers have also been used to enhance and control nucleation, reduce the surface roughness, and improve adhesion of the film to the substrate, the focus of this chapter is on synthesis.
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It is not our intent to provide a general background on lasers or on their principles of operation. There exist several excellent books on the subject [9,10]. Laser operation and characteristics will be discussed only as and when specifically applicable to a particular process. The focus here is on the more nonstandard and novel applications of lasers for diamond growth, especially those that result in highly nonequilibrium environments. In other words, we explore the situations that can best be created by a laser and, in some cases, not by any other means. This chapter is organized as follows. Section II describes synthesis of diamond and diamond-like carbon reportedly resulting from rapid heating and/or cooling due to laser irradiation. Production of diamond-like carbon by ablation of a precursor target is described in Sec. III, followed by a discussion of use of laser CVD for diamond synthesis by photolysis of carbonaceous precursors in Sec. IV. Diamond synthesis by relatively low-power resonance absorption of laser energy (i.e., laser excitation) is discussed in Sec. V. A chapter summary is given in Sec. VI.
II.
RAPID HEATING AND COOLING
A.
Introduction
There are specific instances where laser beams provide an intense source of energy that results in heating of a target gas, solid, or liquid. Depending on the rate at which the incident energy is absorbed and redistributed per unit time, resulting temperatures and pressures can exceed thousands of kelvins and many hundreds of kilobars or more. Referring back to the (equilibrium) phase diagram for carbon (see Fig. 13), it can be seen that such conditions are conducive for conversion of nondiamond carbon to diamond. Even when the output power of a laser beam alone does not provide sufficient intensity or the required heating and compression, the beam can be focused to provide intense heating over a small region. The output beam from a laser is usually at least a few millimeters in diameter. When focused using a lens (of radius b < a) or focusing mirror, the ultimate radius of the beam of radius a can be estimated from physical optics to be [6,9] r0 = f/min{a, b}
(4)
where is the wavelength of laser light and f is the focal length of the lens or optical element used. In practice, however, the ultimate size of a focused laser beam is much larger than that given by Eq. (4) because of multimode emission from the laser and lens aberration. Now, if P is the power of the incident laser beam, then the intensity can be written as
Laser-Assisted and Optical Pumping Techniques
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339
a
I=
Ip(r)2 r dr
0
r 20
When Ip is independent of r, we have I=
P r 20
(5)
We can also write Eq. (5) using Eq. (4) as I = Pa2/2 f 2
(6)
Usually, Ip depends on r in a Gaussian manner but Eqs. (5) and (6) can be easily rewritten for this case. Although Eq. (6) may not be quantitatively accurate for real laser beams for which Ip is a function of r, it does provide the correct trends. For instance, the intensity does scale inversely with the square of the laser wavelength and lens focal length and directly with the square of the initial (i.e., prefocusing) beam size. Thus, for a given laser output power, the intensity increases for shorter wavelength laser beams, smaller focal length lenses, and large initial beam diameters. B.
Gas-Phase Heating
An appropriate example of rapid gas-phase heating triggering reaction and ultimately leading to formation of diamond particles is the experiment of Buerki and Leutwyler [11]. This work combines both pyrolytic and photolytic decomposition of the gas phase. In this experiment, mixtures consisting of ethylene (C2H4), H2, and silane (SiH4) in compositions of 51.5–100%:0– 41.2%:0–7.2% were introduced vertically through a capillary nozzle into a reactor. These precursors were diluted in either N2 or Ar or both. The gas flow was then irradiated horizontally by a focused line-tunable continuouswave (cw) CO2 laser beam, 2–3 mm above the nozzle. Silane (a strong absorber of 10.6-m CO2 laser radiation) was used to elevate the gas temperature and to enhance pyrolysis or chemical decomposition of C2H4. Two processes that cause such gas mixtures to react are pyrolysis and photolysis. The process by which these small molecules decompose upon exposure to infrared radiation from the CO2 laser is known as infrared multiphoton dissociation or IRMPD [7]. Both SiH4 and C2H4 absorb infrared photons via multiphoton processes (of the type discussed in footnote 7 and shown in Fig. 8) triggered by the intense radiation field of the focused laser beam and are subsequently taken to their respective dissociation limits [7]. Inevitably, some of the laser energy is also transferred to the vibrational modes of these
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molecules, resulting in rapid gas heating via vibration-rotation-translation (VRT) and vibration-translation (VT) collisional energy transfer. Thus, some of these molecules are also pyrolyzed and decomposed. The authors report that 1 g of particles black in appearance were produced from C2H4 /N2 /Ar/ H2, C2H4 /N2, C2H4 /Ar, and C2H4 /N2 /H2 mixtures. The particles (on the order of 5–15 nm in size) were characterized by bright-field transmission electron microscopy (TEM) imaging and electron diffraction. Based on this analysis, the authors report finding diamond for the C2H4 /N2 gas mixture and a graphite-diamond mixture in all the other cases. In addition, these authors report that a gas mixture of C2H4 /N2 /SiH4 /H2 yielded diamond particles ⬃120 nm in size in one instance. Although it is encouraging that they appear to be able to distinguish between diamond and graphite through bright-field TEM imaging and electron diffraction analysis, some caution is warranted. Disordered graphite is known to masquerade as diamond, especially in electron diffraction patterns [12,13]. The only way to establish deposits unequivocally as being diamond is through Raman spectroscopy. As an example, the Raman spectrum of a diamond film containing graphite and disordered graphite synthesized using an oxyacetylene flame is shown in Fig. 9. This particular film can be seen to contain diamond (1332 cm⫺1 line), graphite (as can be seen from the ‘‘G’’ peak at 1580 cm⫺1), and disordered graphite
Figure 9 Raman spectrum of a diamond film synthesized using an oxyacetylene flame, displaying peaks identifying diamond and disordered graphite. (From Refs. 15–17.)
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(as evidenced by the presence of the ‘‘D’’ peak at 1343 cm⫺1). Unfortunately, no Raman spectra are presented in Ref. 11. This is understandable given the small size of the particles and the small sample, both of which would yield an extremely small signal (if any) in Raman spectroscopy. In the absence of such evidence, as many additional methods as possible of characterization of the physical properties (such as hardness) and chemical properties (such as reactivity to hydrogen or oxygen when heated, or exposure to acids) are needed. Again, this is difficult when sample sizes are small. Gas-phase reactions with C2H4 as the carbonaceous precursor required about 10 W of power (the residence time was about 1 msec and the pressure was 1 atm) [11]. Based upon the values given in Ref. 11, a characteristic J/carat value (not counting capital cost) can be computed for this case using the optimistic assumption that all of the 1 g of deposit yield is diamond. Reference 11 reports a linear growth rate of 5.7 m/sec, and assuming this is all diamond, we can compute an equivalent spherical volumetric growth rate of (4/3)(5.7 ⫻ 10⫺6)3 m3/sec or 7.8 ⫻ 10⫺16 m3/sec. Given the density of diamond (3515 kg/m3) and that 1 carat = 0.2 g or 2 ⫻ 10⫺4 kg, the growth rate in carats per second is 7.8 ⫻ 10⫺16 ⫻ 3515/2 ⫻ 10⫺4, or 1.4 ⫻ 10⫺8 carats/sec. The power required for this process according to Ref. 11 was 10 W or 10 J/sec, so that the energy cost in joules per carat is 10/1.4 ⫻ 10⫺8, or 7.1 ⫻ 108 J/carat. The value of 7.1 ⫻ 108 J/carat is enormous, and it is important to put it in perspective by comparing with the corresponding energy cost for diamond synthesis using an oxyacetylene flame. For flame synthesis, assuming a linear growth rate of 100 m/hr on a 1 cm2 surface area, we obtain a typical value9 of 8 ⫻ 105 J/carat. Thus, we can see that the energy cost of this particular pyrolysis/photolysis process is three orders of magnitude higher and has no apparent gain in quality of the diamond or in selectivity. It must be further emphasized that laser efficiency and capital cost have not been considered here. The joule per carat value, however, does provide an index for comparing these different processes. A more recent example of extreme gas-phase heating to the point of ionization is the laser-plasma technique developed by Konov and Uglov [14]. In this work, a powerful (2.5 kW) cw CO2 laser beam was focused using an NaCl lens ( f = 7 cm) into a chamber containing a flowing mixture of Xe, H2, and CH4 (in proportions of 300:30:1) at 1 atm. The authors report growth of diamond films 1 cm2 in area at a rate of 30–50 m/hr on water9
Assuming a growth rate of 100 m/hr over a circular spot 1 cm2 in area, the volume growth rate is 2.8 ⫻ 10⫺12 m3/sec or 4.9 ⫻ 10⫺5 carats/sec. Taking 39 W as that power required to raise a total flow of 1.8 ⫻ 10⫺4 kg/sec of an acetylene-oxygen mixture from 25⬚C to the ignition temperature of 305⬚C, the cost is 8 ⫻ 105 J/carat. It must be pointed out that this value is typical of a laboratory-scale oxyacetylene flame.
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cooled tungsten substrates, verified after growth using Raman spectroscopy. This growth rate is comparable to deposition using a laboratory-scale oxyacetylene flame, except that the power cost is higher by about two orders of magnitude. It must be pointed out, however, that this excessive power cost can be lowered if their process is optimized. This work is, however, very similar in nature to synthesis using the oxyacetylene flame except that conditions such as gas type and purity are more controllable. C.
Surface Heating
An alternative approach of using lasers to provide rapid heating has been demonstrated by several researchers [18,19]. In this instance, a stream of carbon black particles was exposed to CO2 and yttrium-aluminum-garnet (YAG) laser irradiation at various power densities and modes of operation (i.e., cw and pulsed). Conversion of carbon black particles to diamond is reported in Ref. 19. Again, no Raman spectra are given, but the authors claim that TEM imaging, x-ray, and electron diffraction data are able to distinguish diamond particles from disordered graphitic phases. Another report appears to confirm this finding based on TEM imaging, electron diffraction, x-ray photoelectron spectroscopy, and measurements of electrical resistivity [20]. In the latter work, a Q-switched Nd:YAG laser was used to irradiate a 200-nm-thick amorphous carbon (so called diamond-like carbon or DLC) film and to convert it to diamond. However, caution must be exercised about the credibility of these reported results given the lack of any Raman spectra. Further, an earlier study on pulsed Nd:YAG laser irradiation of amorphous carbon films actually reports conversion of the film to graphite at intensities of less than about 1 J/cm2 [21]. They confirm the structural change to graphite by Raman spectroscopy. Laser intensities of 4.5 J/cm2 were reportedly used in Ref. 20 and these intensities are not drastically different from those used in Ref. 21. It is also unlikely that the initial amorphous carbon films were very different. This comparison indicates that characterization methods other than Raman spectroscopy (such as TEM imaging and electron diffraction) may not be able to distinguish diamond from disordered graphite reliably. There has been an interesting attempt to use a laser to provide rapid heating of the surface layers of a single-crystal copper specimen, initially saturated with carbon. Because carbon is not soluble in copper, it can be driven out of the material with relative ease by heating. The technique described in Ref. 22 reportedly formed millimeter-size single-crystal diamond ˚ thick and were formed on the singlefilms. These films were about 500 A crystal copper substrate when irradiated with a laser beam. The presence of diamond was reportedly verified using Raman spectroscopy. This work was
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probably based upon the work of Ref. 23, which reported heteroepitaxial diamond growth by ion implantation of high-temperature carbon into a copper substrate. Unfortunately, neither the result reported in Ref. 22 nor that in Ref. 23 was reproduced in a subsequent attempt by Lee et al. [24]. Their approach essentially consisted of ion implantation of carbon at an ion energy of 50 keV and a flux of 1018 atoms/cm2 in a number of crystalline substrates [Cu(100), Cu(110), Cu(111), Ni(100), Ni(110), and single crystal Co]. The samples were then irradiated using separate laser wavelengths of 248 nm (KrF excimer laser with pulse width of 25 nsec), 308 nm (XeCl excimer laser with pulse width of 15 nsec), and 532 nm (frequency-doubled Nd:YAG laser with pulse width of 7 nsec) with power densities in the range of 1–6 J/cm2. The laser irradiation was carried out in vacuum and in the presence of nitrogen. No additional information was reported in their paper about whether or not the beam was focused or the spot size of the beams. Characterization by Raman spectroscopy, Auger spectroscopy, and TEM showed the resulting carbon films to be amorphous or composed of microcrystalline graphite. Lee et al. were therefore unable to reproduce the results of Refs. 22 and 23. Reference 24 appears to have reproduced all of the conditions reported in Ref. 22, but for the longest pulse width and perhaps the rise time of the laser pulse. In the experiments of Ref. 24, the longest pulse width was 25 nsec whereas the longest pulse width in Ref. 22 was 45 nsec. Because the pulse width affects the heat flux and total heat input into the surface, this could have affected the outcome of the experiments, with the rise time of the laser pulse being an important parameter. Figure 10 shows how two pulses can have different rise times while having the same total or integrated energy content. As we are dealing with pulse widths on the order of tens of nanoseconds, specific details such as rise time of the pulse and maximum pulse amplitude could be important in reproducing the experiments reported in Refs. 22–24. In other words, even if the heat addition over the duration of the laser pulse were the same in Refs. 22 and 24, the
Figure 10 Schematic showing two laser pulses having the same total or integrated energy while having different peak powers and rise times.
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rate at which the heat is added may have been important. Unfortunately, pulse shapes are not reported in either Ref. 22 or Ref. 24. In addition, it would be interesting to repeat these experiments in the presence of atomic hydrogen, as this could help stabilize a diamond surface in the act of formation. It is instructive to explore the heating depth and evolving surface temperature distribution when a material is exposed to laser irradiation. Laser radiation can be reflected, transmitted, or absorbed at a surface and at depths below the surface. Of these, only the absorbed energy affects the temperature or average energy of the solid material. The absorbed radiation may be viewed as the result of interactions between the incident photons and electrons and photons and phonons10 in the material. In dielectric materials, phonon-phonon interactions are dominant whereas in metals photon-electron followed by electron-phonon interactions are important. The process of heat transport within the material therefore establishes itself (in the classical sense of heat conduction) only after a certain time has elapsed. This initial relaxation time, , is the characteristic time required for photon-phonon or photon-electron and electron-phonon interactions (i.e., collisions) to exchange energy and establish local thermal equilibrium. In other words, the relaxation time is the time that must elapse before a local lattice temperature or local electron temperature can be defined. In metals, the characteristic times for phonon-phonon, photon-electron, and electron-phonon interactions are on the order of picoseconds, and in dielectrics the phononphonon relaxation time is on the order of nanoseconds to picoseconds. Therefore, if laser pulse widths are shorter than nanoseconds, the microscale or short-transient energy transfer must be considered, and for subnanosecond time scales it is likely that the temperature field is discontinuous [25]. During these short times, propagation of thermal energy within a material often follows a hyperbolic or wave equation unlike the usual parabolic equation describing the diffusion of heat.11 This means that a temperature distribution can propagate as a pulse or wave through the material, as shown in Fig. 11, with transient peak temperatures being much higher than normally predicted by the heat diffusion equation. Values of the quantities q (relaxation time 10
Phonons are the quantum particles corresponding to the discrete or quantized vibrational states in the lattice of a solid. Typical lattice frequencies are on the order of 1013 Hz and energies are on the order of E = h, where h is Planck’s constant. Phonons propagate at the speed of sound, which for most solid materials at room temperature is on the order of 10 to 100 km/sec. 11 Although the hyperbolic form of the heat equation was first pointed out by Maxwell in the nineteenth century [26], it remained obscure until Cattaneo revived the idea in 1948 and again in 1958 [27].
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Figure 11 Plot showing the distribution of nondimensionalized temperature (x, t) = (T(x, t) ⫺ T0)/T0 versus nondimensionalized distance ␦ = x/2兹␣q, at a nondimensional time  = t/2q. The parameter B = T/2q, where T is the characteristic time for a meaningful temperature to be defined at a point or the phase lag of the temperature gradient, q is the characteristic time for a meaningful temperature gradient to be defined at a point or the phase lag of the heat flux vector, ␣ is the thermal diffusivity, and T0 is the initial temperature at t = 0. The limit when B = 0 corresponds to what is known as a thermal wave, i.e., one in which a temperature pulse with an overshoot propagates as a wave through the material medium. (From Ref. 25.)
for the temperature gradient to be established) and T (relaxation time for a temperature to be defined) inferred from measurements are 0.4348 and 70.833 psec, respectively, for metallic copper [25]. However, the corresponding values for oxides are larger, and it must be borne in mind that copper will have an oxide layer when exposed to air. Such differences can be important in explaining the different results obtained in Refs. 22–24. The effects of heating at short time scales (i.e., on the order of nsec or smaller) can be considerably different from behavior at long time scales (or times greater than a nanosecond), and this in turn can alter the diffusion of carbon from the bulk material to the surface.
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Following the initial transient, the temperature response of a solid subjected to laser irradiation can be calculated for times on the order of tens of nanoseconds and greater. This is done by ignoring a part of the fast rising portion at the beginning of the laser pulse. In this instance, the heat flow in the material can be determined for most of the pulse duration of tens of nanoseconds using classical (Fourier-type) heat conduction. For a semi-infinite solid receiving a heat flux over a small, prescribed region on its surface, the transient temperature distribution is [28] T(x, t) ⫺ T(x, t = 0) = ⫺
q⬙x 0 erfc kT
1/2 2q⬙( 0 ␣Tt/) exp kT
冉 冊 x
2兹␣Tt
冉 冊 ⫺
x2 4␣Tt
(7)
where q⬙0 is the heat flux at the surface due to the absorbed laser energy, ␣T is the thermal diffusivity, kT is the thermal conductivity, x is the spatial coordinate, t is time, and erfc is the complementary error function. T(x, t = 0) is the temperature distribution following the fast transient portion of the laser pulse, after which time classical heat conduction occurs. The q⬙0 in Eq. (7) is a constant heat flux input into the surface. In reality, q⬙0 can be time dependent and varying even in the regime of classical heat conduction for the two pulse shapes shown in Fig. 10, so that Eq. (7) must be rederived in such an instance. In addition, the boundary condition (which yields the instantaneous value of q⬙) 0 must include a heat balance whereby heat is removed from the surface by radiation and convection. Hence, unrealistically high surface temperatures will be predicted by Eq. (7) if applied under conditions where radiative cooling of the surface is important. Although the situation described by Eq. (7) is not exactly that encountered in the experiments [22–24], it can nevertheless provide an indication of the temperature levels existing within the material during the period of laser irradiation. For illustrative purposes, temperature profiles given by Eq. (7) are plotted at various times in Fig. 12 for a laser intensity of 5 ⫻ 108 W/cm2, representative of the conditions existing in the experiments of Ref. 24. Because the rate of diffusion of carbon within the bulk Ni or Cu substrate material is temperature dependent, specifically varying as T 1/2, carbon initially near the surface would be driven out, aggregate, and condense to form a film. Although this would be expected to affect the growth rate of the carbon film at the surface, it is unclear what the effect of high local transient temperatures would be on the type of carbon aggregates thus produced. The nature of the resulting film could well be dependent on the rise time of the laser pulse.
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Figure 12 Temperature distributions calculated using Eq. (7) at various instants of time are shown here for an absorbed laser intensity of 5 ⫻ 108 W/cm2 for copper, representative of conditions in the experiments of Ref. 24.
Perhaps the most dramatic and as yet unsubstantiated claims have been the reports of diamond growth on a variety of substrates using CO2 as the carbonaceous precursor, with one or more laser sources (a CO2 laser, an excimer laser, and an Nd:YAG laser) [29–32]. It is important to point out that these claims were reported initially in the popular press, thereby circumventing the usual scientific peer review process. However, they are worth examining in view of what already exists in the published scientific literature and what is discussed in Secs. IV and V of this chapter. We will try to analyze this process, henceforth referred to as the QQC process (named after the industrial group that purportedly invented it). According to initial press releases, in the QQC process CO2 and N2 (as a shielding gas) are introduced in a plane stagnation flow over the substrate at an initial pressure of 1 atm. The process then consists of simultaneously irradiating a target surface to be coated, with simultaneous focused
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beams from a CO2 laser, excimer laser (of undisclosed wavelength), and an Nd:YAG laser, which then reportedly produces diamond on a variety of substrates including stainless steel, cast iron, and a variety of metals and ceramics [29–32]. What is more remarkable is the claim of growth rates on the order of 1 m/sec (specifically, a 45-m diamond coating is reportedly produced over a 1.6 cm2 area within 45 sec). Although few details are provided in Refs. 29–32, it is still possible to explore the feasibility of this approach. We now analyze the QQC process to explore whether or not the conditions might even be conceivably favorable for diamond growth. The excimer lasers and YAG lasers used in the QQC process are capable of producing extremely short pulses on the order of 10 nsec, so that the peak power can be on the order of 20 MW12 or more. This implies a peak intensity of 20 MW/mm2 or 2 GW/cm2 for a spot size or focal area of 1 mm2. An estimate of the gas temperature can then be made if we assume that the laser heating is balanced by radiative loss from the gas. This yields an approximate value of 13,000 K for the gas temperature of the resulting thermal plasma. Ionization and dissociation processes will tend to lower this estimated temperature so that realistic temperatures are probably in the range of thousands of kelvins. In the reported experiments, the target surface is exposed to a high intensity of infrared and ultraviolet energy. This must result in flash evaporation, in a manner similar to what occurs in laser shock processing or explosive welding [33–36]. The flash evaporation of the substrate material can result in an implosive (compressive) force at the material surface in the gigapascal range, depending on the pulse energy.13 This effect is put to good use in laser shock processing to improve the mechanical properties of metallic surfaces of dense or porous materials and in explosive welding [33–36]. The resulting gas mixture composed of vaporized substrate material, CO2, and N2 is rapidly heated to temperatures high enough to ionize the constituents. Such an ionized, electrically conducting gas is called a plasma. Accompanying this rapid heating is the production of extreme pressures. For a Gaussian-shaped laser pulse, assuming an absorption coefficient of 0.1 (typical for plasmas in this range of temperature and pressure), Ref. 35 gives the following relation between the maximum pressure in kbar and the absorbed intensity expressed in GW/cm2:
12
This is obtained assuming a peak energy of 200 mJ and a pulse width of 10 nsec. ˙ mass per unit time. This is typically milliA surface undergoing flash evaporation loses m ˙ on the order of 102 kg/ grams of material over the pulse duration of 10 nsec, yielding an m sec. Assuming a shock velocity on the order of the speed of sound at the temperature of 13,000 K, or O(103) m/sec, this yields a force of magnitude 105 N, or a pressure on the order of 1011 Pa (100 GPa, or 1000 kbar) on a spot 1 mm2 in area.
13
Laser-Assisted and Optical Pumping Techniques
Pmax = BI 1/2
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(8)
where B is a constant with a value on the order of 10. It is interesting to note that this expression yields a maximum pressure that is independent of laser pulse duration and wavelength and one that is dependent only on the laser intensity at the focal point. Therefore, it is evident that the local processing conditions (i.e., in the gas and on the surface) achieved by such powerful lasers are on the order of GPa of pressure and temperatures on the order of thousands of kelvins. Shown in Fig. 13 and Fig. 14 are the phase diagrams for carbon in the absence of a catalyst and in the presence of catalytic metals such as nickel and iron, respectively. It can be seen that the pressures and temperatures shown in Figs. 13 and 14 match those expected in the QQC process. Thus, the QQC process appears capable of producing conditions locally comparable to those encountered in the well-known hightemperature, high-pressure (HTHP) process used for commercial diamond growth [37]. This places the QQC process within the realm of possibilities. However, the technique must be reproduced and independently verified before its claims such as growth rates of 1 m/sec over an area of 1.6 cm2 can be substantiated. According to the analysis presented here, such growth rates may be possible on surfaces approximately 1 mm2 in area.
Figure 13 Ref. 38.)
Phase diagram for carbon without the presence of any catalyst. (From
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Figure 14 Phase diagram for carbon in the presence of Ni catalyst (left) and Fe catalyst (right) at a pressure of 5.7 GPa. M, solid metal; L, liquid; G, graphite; D, diamond; C, carbide. (From Ref. 38.)
We can estimate the cost in terms of joules per carat for the QQC process, assuming that the growth rate of 1 m/sec over a 1.6 cm2 area is indeed realizable. Using a pulse energy of 200 mJ/sec and volume growth rate of 1.6 ⫻ 10⫺10 m3/sec, we can calculate an energy cost of 71 J/carat, which is energetically much cheaper than diamond synthesis by one of the faster CVD processes, flame synthesis. D.
Heating of Surfaces Immersed in Liquids
The growth of some natural diamond (if not all of the larger crystals found in nature) is believed to occur in a state comprising solids immersed in
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liquids at extreme pressures and temperatures [38]. Nevertheless, synthesis of diamond by various deposition techniques has until recently largely focused on the gas phase alone. There have been a few attempts to grow diamond from a liquid phase or from solids submerged in the liquid phase [39–43]. Of these, Refs. 40–42 appear to be most intriguing in terms of potential application of lasers for diamond growth. In Cherian’s experiments [40–42], evidence for dissolution of diamond in a nickel crucible containing liquid NaOH at 1100 K and at atmospheric pressure is given, and more surprisingly, recrystallization of diamond back onto diamond surfaces upon ‘‘rapid’’ cool-down is also observed. As reported in Ref. 40, diamond crystallites consisting of chips and cleavages were dissolved in molten NaOH in a nickel crucible maintained at 1100 K in an oven. The crucible and its contents were then rapidly cooled to room temperature by removing them from the oven and placing in contact with a room-temperature metal surface or by cooling the crucible in room-temperature water. A cooling rate of ⬃1000⬚C/min was reportedly required for recrystallization to occur. Figures 15 and 16 show scanning electron microscopy (SEM) images of overgrowths that show the recrystallized diamond. Verification that the overgrowths are indeed diamond was carried out using several different characterization techniques, including microprobe Raman spectroscopy [40]. Reference [40] reports that careful experiments varying the quenching rate showed that the faster quenching rates yielded incomplete growth in the direction perpendicular to the substrate. Further, a metallic film of Ni was
Figures 15 and 16 SEM images of overgrowths that show the recrystallized diamond upon quenching an NaOH solution containing dissolved diamond particles in a nickel crucible. (From Ref. 41.)
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found on the overgrowths of diamond, suggesting the role of an Ni-C complex as an intermediate facilitating the growth of diamond. Evidently, some of the carbon in the initial diamond seed material dissolved into the liquid NaOH, formed complexes with the dissolved Ni (from the walls of the Ni crucible), and upon cool-down recrystallized as diamond on remnant diamond seed surfaces. More insight into mechanisms could have been gained if isotopic labeling had been used. For instance, 13C diamond particles (used in the dissolution part) and 12C diamond seeds (used in the recrystallization part) may have been helpful in identifying the exact sources and targets of dissolution and recrystallization, respectively. Alternatively, a carbonaceous source other than diamond (graphite, for instance) could have been used in the dissolution phase of the experiment, while a diamond or silicon seed particle (thoroughly characterized prior to use) could have been introduced during the growth/recrystallization or cool-down phase of the experiment. More recently, Zhao et al. [43,44] have demonstrated that diamond can be grown on diamond seeds from (Ni)-C-H2O mixtures at a pressure of 1.4 kbar and temperature of 800⬚C.14 Experiments typically used mixtures of 3 wt% powdered nickel, 95 wt% glassy carbon, and 2 wt% 0.25 m sized diamond in water (50–100 wt% of the glassy carbon) in a pressurized and heated vessel. Run durations were typically 50–100 hr, the shorter durations yielding diamond particles several microns in size. The presence of nickel was reportedly crucial in obtaining larger aggregates of particles with sizes ranging from tens of microns to 100 m. The typical particle morphology observed in these experiments is shown in Fig. 17 and can be seen to resemble the morphology of diamond particles produced by General Electric’s commercialized high-pressure, high-temperature process. The use of lasers in such liquid-phase experiments may be awkward, although specific wavelengths do penetrate certain liquids. For example, the green line of the argon ion laser can be transmitted through water. Therefore, it may be possible to use lasers to transmit through a liquid and trigger heterogeneous reactions at a liquid-solid interface. However, this can be exploited only after growth mechanisms are clearly understood. Another 14
Laboratory experiments (see H. C. Noltimier and V. V. Subramaniam, Laboratory chemical vapor deposition [CVD] experiments applied to the geological formation and age of natural diamond, Poster Paper BTH-1 presented at Session 152 [Experimental Petrology] at the 1995 Geological Society of America Annual Meeting, New Orleans, November 9, 1995) also suggest the beneficial role of water vapor in accelerating growth of diamond by hot-filament CVD on minerals such as almandine and olivine, both found as impurities in natural diamond. These minerals can act as nuclei for diamond growth in C-H-O systems (and even at the low pressures of 30 torr typical in HFCVD). The authors also give compelling arguments indicating that most natural diamond may form during transit in the diamond diatremes instead of forming at depth, as commonly thought.
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Figure 17 Typical particle morphology observed in the experiments of Ref. 43. Note the resemblance to the morphology of diamond particles produced by General Electric’s commercialized high-pressure, high-temperature process.
means of possibly initiating diamond growth from a laser-excited liquid is discussed in Sec. V.C.
E.
Summary
Use of intense laser irradiation for pyrolysis (and subsequent reaction) of the gas phase and rapid heating of surfaces has been discussed in this section. Where rate of heating is necessary and crucial, lasers may be capable of providing a unique avenue for diamond growth. However, capital and maintenance costs remain prohibitive, especially in the energy cost per carat. Some reported techniques such as the QQC process, if reproduced, appear competitive with other existing techniques based upon reported growth rates and should be examined more carefully. Lasers have also been used to irradiate powders of carbon black and convert some of these into diamond. A relatively unexplored area is the use of lasers to drive heterogeneous reactions on surfaces immersed in liquids. Use of lasers to excite liquids themselves and drive liquid-phase reactions is discussed in Sec. V.
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ABLATION
The use of lasers to produce rapid heating is well established. When this rapid heating of solid surfaces leads to vaporization of the material at the surface, this is known as ablation. The resulting vaporized material often constitutes an activated gaseous phase that can be subsequently reacted and modified to yield a product or can condense onto another surface to produce a solid layer with different microstructural properties. This process, also known as laser physical vapor deposition, was shown schematically in Fig. 7. Intense laser energy (usually focused) is directed toward a target material, causing it to vaporize. The resulting vapor or plume is a gas comprising atoms and molecules from the target and is often ionized. This technique is often used to produce thin films on substrates by directing the electrically conducting plume toward a biased substrate. LPVD or ablation has been used for the deposition of thin films of different materials. Of particular relevance is the use of such an ablative technique for the deposition of diamond-like or amorphous carbon films [45–54]. In these experiments, graphite is used as a target (except in the case of Ref. 54, where Lucite targets were also used) and different highenergy lasers such as excimer lasers or Nd:YAG lasers or CO2 lasers are used. Intensities on the order of 108 W/cm2 are usually required. It is well established that LPVD is capable of synthesizing diamondlike carbon (DLC) or amorphous carbon films from ablated graphite targets. The term amorphous carbon refers to a material with properties distinct from those of graphite but lacking the long-range crystalline order characteristic of diamond. DLC films are known for their hardness approaching that of diamond. Many of the properties of amorphous carbon materials can be accounted for by the structural model of graphitic islands interlinked by small percentages of diamond-like sp 3 bonds [55]. Davanloo et al. [55] report the growth of ‘‘amorphic diamond’’ from a laser plasma source. These authors provide sufficient evidence for differences between their material and DLC (such as structural size on the order of 100–200 angstroms and band gap of 1.0 eV) to warrant distinguishing between their material and other established DLC properties. Such small crystallite dimensions usually pose great difficulty in acquiring Raman spectra so that this characterization technique is ineffective in unequivocally indicating the presence of diamond in the present case. Their technique is a variant of LPVD. An Nd:YAG laser (Q-switched repetition rate of 10 Hz, 15 nsec pulse width) was used to produce a peak intensity of 5 ⫻ 1011 W/cm2 at a graphite target, in conjunction with an electric discharge of 10 A maintained between the target and a rod electrode placed nearby. The special nature of this ‘‘amorphic diamond’’ film thus produced on a germanium window is evident in Fig.
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Figure 18 Transmittance curves for an amorphic diamond–coated Ge window (25 mm diameter) compared with a bare Ge window. The film was reportedly 0.3 m thick with an index of refraction of 2.1. The authors point out the fact that absorption from CH bands at 2940 cm⫺1 is absent. (From Ref. 56.)
18. As can be seen from Fig. 18, the transmittance is higher for the amorphic diamond–coated Ge window compared with the bare Ge window and shows complete absence of the C — H absorption band at 3.4 m. Reference 55 further reports a hardness value of 13 GPa, compared with the value of 77 GPa (at a temperature of 290 K) for diamond [37]. The amorphic diamond of Davanloo et al. was thus clearly subjected to numerous such tests, which enabled them to confirm that their films indeed exhibited diamond-like properties. Although LPVD and plasma-assisted LPVD have been successful in producing DLC and amorphic diamond films, there have only been a limited number of reports regarding growth of diamond using this technique [56,57]. Latsch and Hiraoka have shown that diamond can be produced by LPVD from targets of polymethyl methacrylate (PMMA) when a reactive mixture of oxygen and hydrogen (ratio of concentrations not reported) is present at pressures between 0.5 and 0.9 torr. The authors report the following critical parameters: target to substrate distance = 3–6 cm, substrate temperature = 600⬚C, ArF excimer laser (193 nm), fluence = 190 mJ/cm2, and substrate
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bias = ⫺50 V [56]. The authors do not provide any information regarding any substrate pretreatment and report that ‘‘crystals started to appear after 10 minutes and thicker films were obtained for longer exposure times.’’ However, no SEM images are shown of the ‘‘films,’’ only of the isolated particles (see Fig. 19). Further, microprobe Raman spectra of the deposits (shown here in Fig. 20) indicate the presence of both ‘‘D’’ (at 1330 cm⫺1) and ‘‘G’’ (1597 cm⫺1) peaks. These peaks are usually indicative of graphite. Although the ‘‘D’’ peak usually occurs around 1343 cm⫺1, it can shift in some samples to the location where the diamond peak occurs. Therefore, caution must be exercised in interpreting Raman spectra especially when a broad peak near 1332 cm⫺1 and the ‘‘G’’ peak of graphite are simultaneously present. As an illustrative example, a Raman spectrum of a carbonaceous deposit produced from an oxyacetylene flame is shown in Fig. 21 [17]. The same film, exposed to atomic hydrogen (produced by a hot filament), removed the carbonaceous deposit, as can be seen from the Raman spectrum shown in Fig. 22 [17]. Given the uncertainty introduced by this fact, it is important to expose a film (or crystals) suspected to be diamond to H atoms at temperatures near 800⬚C in order to be sure that the deposit is indeed diamond. Subsequently, Polo et al [57] used ArF and KrF excimer laser ablation of graphite targets in hydrogen at 0.76 torr and reported growth of diamond particles 1–20 m in size. The deposits were verified as diamond by Raman spectroscopy. The ablative method can be subject to problems of reproducibility. The composition of the target is often a critical parameter, as are the substrate temperature, position of the substrate relative to the target, substrate bias,
Figure 19
SEM image of isolated particles produced by LPVD. (From Ref. 56.)
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Figure 20 Microprobe Raman spectrum of the deposits from Ref. 56 showing the presence of both ‘‘D’’ (at 1330 cm⫺1) and ‘‘G’’ (1597 cm⫺1) peaks. This is indicative of disordered graphite rather than diamond.
Figure 21 Raman spectrum of a carbonaceous deposit produced from a flame. (From Ref. 17.)
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Figure 22 The film whose Raman spectrum is shown in Fig. 21 was exposed to atomic hydrogen (produced by a hot filament) [17]. This figure shows the Raman spectrum after exposure to atomic hydrogen. Note the disappearance of the feature at 1334 cm⫺1, which clearly demonstrates that the original film was not diamond.
and laser pulse energy. Of these, variability in the composition of the target has the largest influence on the resulting film or deposits. In some cases, ablation of the target can be accompanied by gross ejection of particulate solid material from the target that is then incorporated in a growing film. Based upon the scant information presently available regarding growth rates in LPVD, it is not possible to estimate the energy cost per carat for this process. Further research in this area is warranted, especially with regard to use of different combinations of target material and reactive gases.
IV.
PHOTOLYSIS
The process of photolysis, or the use of energetic photons to fragment a polyatomic molecule, was introduced in Sec. I. Usually, such processes require the presence of intense laser irradiation to initiate interaction between a molecule and one or more photons. As explained in Sec. I, this multiphoton interaction between molecule and photons may be sequential (i.e., the molecule may sequentially absorb each photon) or simultaneous, ultimately lead-
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ing to fragmenting of the molecule into smaller molecular or atomic fragments. This generally falls within the area of photochemistry [4,7,8]. The wavelength range typically used in photochemistry varies from the vacuum UV (below 200 nm) to the microwave region (1–10 cm). This range is chosen because of the typical absorption characteristics of molecules. Most small (three or four atom) polyatomic molecules have bond dissociation energies on the order of 1–10 eV. This corresponds to photon wavelengths in the UV-visible range. Light in this wavelength range is usually absorbed via electronic transitions in molecules. Wavelengths in the IR usually result in vibrational excitation because molecular vibrational states differ by energies on the order of 0.1 eV or less. Radiation in the far IR to microwave regions (10 m to 10 cm) usually corresponds to excitation of the rotational energies of molecules. The wavelength range applicable to photochemistry is further restricted by the availability of materials for optical elements such as windows, lenses, and mirrors. Typical wavelength ranges and other properties for common optical materials are listed in Table 2. Table 3 lists some useful conversion factors because different units are used for convenience. For instance, when dealing with UV or visible light, wavelength units in nanometers are typically used. When dealing with radiation in the near IR to the far IR, units of microns (m) are used for wavelength and wavenumbers or cm⫺1 are used for frequency. In the microwave regions, wavelength units of millimeters or centimeters are used. It is important to delineate the difference between reactions occurring in photochemistry, thermally induced reactions (i.e., those caused by heating), and reactions initiated by extremely short-wavelength radiation such as x-rays or gamma rays. In thermal reactions (such as pyrolysis), molecules are dissociated or reacted from their ground electronic manifolds by vibrational excitation usually fed by energy from the external modes (i.e., translation and rotation). In contrast, photochemical reactions typically involve electronically excited states. In reactions initiated by x-rays and gamma rays, the route is often via ionization. In other words, molecules are ionized (i.e., charged by stripping them of one or more electrons) and subsequently react with other ions or neutral molecules. The difference between thermally induced reactions and photolytic reactions is apparent in the following illustration provided in Ref. 4. Consider the thermal decomposition of ethane: C2H6 ⫹ M → CH3 ⫹ CH3 ⫹ M where M is any other collision partner. Here, translational energy is transferred from M to C2H6 by impact, resulting in an intermediate15 fol15
The intermediate (C2H6)† is typically vibrationally excited and is known as a transition state. In contrast, the intermediate in photolytic reactions is usually an electronically excited state.
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Table 2
Typical Characteristics of Optical Materials
Material
Source: Ref. 59.
0.15–5.5 0.5–13 0.12–10 1.8–22 0.12–7 0.11–8 0.21–37 0.18–25 0.16–15 0.27–42 1.2–15 0.12–11 0.395–5.0 0.42–5.3
1.768 1.475 at 0.546 m 1.424 at 2 m 4.116 at 2 m 1.379 at 2 m 1.378 1.538 at 2 m 1.475 at 2 m 1.39 at 0.257 m 1.631 at 2 m 3.458 at 2 m 1.439 at 0.546 m 2.409 2.613–2.909
Insoluble
2200
3.97
2053
0.12 1.6 ⫻ 10⫺3 Insoluble 0.27 7.6 ⫻ 10⫺3 53.48 34.7 92.3 127.5 Insoluble 0.011 Insoluble Insoluble
82 158 780 102 415 7 9
4.89 3.18 5.323 2.635 3.15 2.75 1.984 2.48 3.12 2.33 4.24 5.175 4.26
1280 1360 937.4 842 1266 730 776 846 682 1410 1450 2080 1825
5 1150 130 595 700
Cleavage plane
Thermal expansion (⫻10⫺6 per ⬚C) 4.5
(111) (111) (100) Poor a or c (100) (100) (100) (100)
18.4 6.1 35 9.2–13.4 38.4 36.6 32 43 3
(111) 10.6 C-axis: 9.943
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Al2O3 (sapphire) BaF2 CaF2 Ge LiF MgF2 KBr KCl KF KI Si SrF2 SrTiO3 TiO2 (rutile)
Transmission Water solubility Melting range (g/100 mL of Knoop Density point (m) Index of refraction H2O) hardness (g/cm3) (⬚C)
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Conversion Factors for Energy
1 eV
1 cm⫺1
(angstroms)
1.6021 ⫻ 10⫺12 ergs/molecule 8065.73 cm⫺1 23.061 kcal/mole 96.487 kJ/mole 2.859 cal/mole 11.962 J/mole 1.9863 ⫻ 10⫺16 ergs/molecule 1.2398 ⫻ 10⫺4 eV/molecule 285915/ kcal/mole 12398/ eV/molecule
lowed by decomposition into CH3 molecules. In contrast, the primary photochemical channel for photolytic decomposition of ethane is C2H6 ⫹ h → C2H4 ⫹ H2 proceeding via absorption of a photon by C2H6, leading to an electronically excited intermediate. As can be seen from this illustration, the products resulting from thermal reactions and photolytic reactions can be very different, showing how different products arise from different transition states. A photon absorbed by a molecule does not necessarily lead to a chemical change. The molecule certainly becomes excited upon absorbing the photon, but there are other loss processes that inhibit chemical change. For instance, the excited molecule could simply radiate the photon away by spontaneous emission. The molecule could also suffer a collision with another molecule while in this excited state. This may result in energy transfer to (into either internal or external modes of) the collision partner. Such processes are considered as losses because they do not result in dissociation of the target molecule or in chemical reaction, which is the primary goal of photolytic processes. In that sense, these loss mechanisms lead to inefficiencies in the photolytic process. The efficiency of a given photolytic process is qualitatively accounted for by the quantum yield, q. It is defined as [60,61] ⌽q =
number of molecules undergoing the process number of photons absorbed by the system
(9)
In photochemistry, 1 mole (i.e., Avogadro’s number of 6.022 ⫻ 1023) of photons is called an Einstein. Typically, an overall reaction pathway may consist of several primary or elementary pathways and a quantum yield can be defined for each primary channel i:
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⌽i =
number of excited molecules that proceed via elementary channel i number of photons absorbed by the system (10)
If moderate or low light intensities are involved, the probability that an excited molecule will absorb more than one photon during its short lifetime is small. For these conditions,
冘
⌽i = 1
(11)
i
where i represents the ith primary photochemical pathway. Equation (11) is referred to as the Stark-Einstein law and holds only for low or moderate light intensities. Further, though ⌽i ⱕ 1, the overall quantum yield for a product, ⌽oq, can be much larger than 1 as is the case for chain reactions. The earliest attempt to use a laser to photodissociate a carbonaceous precursor and produce diamond met with failure [13]. Although success was reported initially [12,13], subsequent characterization of the growth material showed the deposit to be heat-treated carbon black [13]. At the time, both CH4 (methane) and C2H2 (acetylene) were suspected of being growth species and Kitahama et al. used 193-nm radiation from an ArF excimer laser to photodissociate C2H2 diluted (from 1 to 30%) with hydrogen. The pressure was kept in the range of 25–30 torr, and substrate temperatures in the range 400–800⬚C were explored. After growth durations of 2 hr during which time C2H2 was photolyzed using 193-nm pulses (100 Hz repetition rate and 50 mJ/pulse), growth of particles (later shown to be nondiamond carbon) was observed along scratches induced by substrate pretreatment.16 The authors do not explain why higher hydrogen concentrations were not used in their mixtures, especially given the fact that the photochemistry of acetylene at 193 nm is well known [62]. In fact, the photolysis products of pure C2H2 have been reported to yield diacetylene, ethylene, hydrogen, and a polymer [62]. It is quite possible then that higher hydrogen concentrations or presence of oxygen in the precursor gases would have yielded diamond growth in these early experiments. The initial claims of diamond growth by photolysis of acetylene were subsequently retracted [13]. Despite the initial setback for photolysis, Goto et al. [63] succeeded in producing diamond from a gaseous mixture of hydrogen flowing at 1000 sccm and carbon tetrachloride (CCl4) flowing at 1–10 sccm and irradiated by a focused ArF excimer laser beam. The pressure ranged from 6 to 50 torr, the repetition rate was 100 Hz, the pulse energy was 80–200 mJ, and the average power was 5–12 W. The beam was focused using an F:300 16
Substrates were pretreated as is common with other growth techniques, in this case using no. 1800 diamond powder to enhance nucleation and accelerate growth.
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cylindrical lens, and the authors report that configurations where the laser irradiated the substrate and was parallel to the substrate, were separately explored. The latter arrangement is shown schematically in Fig. 23. The authors report the formation of an amorphous carbon film after deposition durations ranging from 30 min to 3 hr. However, when the hydrogen stream was predissociated using a microwave discharge or a resistively heated hot filament, the authors report growth of a 1- to 3 m-thick diamond film on silicon substrates heated to 450⬚C. The films were characterized by Raman spectroscopy, which revealed a discernible peak at 1333 cm⫺1 confirming the presence of diamond. The crystal structure was further confirmed by electron diffraction. This was the first reported instance of diamond growth using a laser-based process and one of the first to demonstrate that substrate temperatures did not need to be as high as 800⬚C. The authors report on the importance of atomic hydrogen and methyl radicals for diamond growth. This idea has received much support in reviews of diamond growth mechanisms [64–66]. Following the success of Goto et al., Tyndall and Hacker [67] reported on photolysis of over 20 compounds of organic precursors using focused (cylindrical lens with a focal length of 100 mm) KrF excimer laser radiation. Laser fluences of 150 mJ/pulse/cm2 were reportedly used along with substrate temperatures between 25 and 400⬚C, while chamber pressures were maintained between 1 and 5 torr. The authors report that temperatures in excess of 400⬚C resulted in dramatically decreased growth rates. The organic precursors were admitted into the growth chamber by passing a carrier gas (reportedly argon, helium, or hydrogen) over the reservoir containing the organic precursor. Of the organic precursors examined, the authors report that aliphatic carboxylic acids such as CH3CO2H (acetic acid) and
Figure 23 Schematic of experimental arrangement used in Ref. 63 to produce diamond films from photolysis of CCl4/H2 mixtures.
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H2C(CO2H)2 (malonic acid) yielded films containing diamond and disordered graphite after durations of 1.5 hr. The malonic acid had to be heated to 50– 70⬚C to provide sufficient vapor pressure to yield reasonable growth rates, but acetic acid could be used at room temperature. The authors do not specify which carrier gas or gases were used with these precursors. The Raman spectra of films obtained from these organic precursors display broad features at 1335 cm⫺1, which the authors claim to be diamond. However, this should be taken as insufficient verification, based upon the illustrative spectra shown in Figs. 21 and 22. As discussed earlier, this film shows the presence of a feature at 1332 cm⫺1, but this feature disappeared when the film was treated in atomic hydrogen (from a heated tungsten filament) after growth. Diamond films are usually quite resistant to etching by atomic hydrogen, and such disappearance of the well-known diamond feature in the Raman spectrum is not observed in true diamond films. On the other hand, graphitic films or films with disordered graphitic carbon show significant changes in their Raman spectra after postgrowth treatment in atomic hydrogen. Therefore, the spectra reported in Ref. 67 cannot be taken as irrefutable evidence of the presence of diamond. A postgrowth etch treatment of the film would have served as another means of verifying their claim. However, the authors of Ref. 67 do report on additional characterization of their deposits, such as electrical resistivity measurements and characterization by Auger spectroscopy. The electrical resistivity of these samples was greater than 109 ohm-cm (closer to the value for diamond, i.e., 1016 ohm-cm, than for graphite, i.e., 10⫺2 ohm-cm). Such additional characterization of the deposits is encouraging and bolsters the claims reported in Ref. 67. No SEM images of the deposits are shown, and therefore it is not clear whether the yield is a film covering a certain region of the substrate or consists of particles distributed over the substrate. In some instances, the choice of the carbonaceous precursor dictates the choice of growth technique. This is true in the case of growth of diamond from mixtures containing CO [68–70]. Use of CO as the carbonaceous precursor for diamond growth is of particular interest when isotopically pure diamond is desired. Diamond synthesized from a hydrocarbon source such as methane, but with a 13C content of 0.1% or less, is much more thermally conductive than diamond grown from CH4 with the natural abundance of 13 C of 1.1% [71]. Now, isotopically pure 12C and 13C are commercially produced as CO, and more processing steps are required to convert them to isotopically pure hydrocarbon sources. Consequently, 12CO will be cheaper than using 12CH4 in such an instance. Furthermore, because CO cannot be used in some conventional diamond growth systems such as HFCVD or oxyacetylene flame synthesis, growth options using CO as the precursor are
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limited. Consequently, plasma CVD or laser methods must be used when CO is the precursor. Growth of particles of diamond on seeded unheated monocrystalline silicon [(001) and (111)] substrates has since been reported from CO/H2 gas mixtures by photolysis of CO [68–70]. The photolysis of CO was accomplished using focused ArF excimer laser irradiation (the reader is referred to Ref. 70 for experimental details). Several orientations of the laser relative to the substrate were explored, as shown in Fig. 24, while the gas flow was directed perpendicular to the laser beam and toward the substrate. A CO flow of 0.7 sccm and H2 flow of 99.3 sccm were used, and the chamber pressure was maintained at 8 torr. Prior to growth, the substrates were thoroughly characterized using SEM imaging and energy-dispersive x-ray spectroscopic (EDS) analysis to ensure that no diamond particles remaining from the abrasive pretreatment used were present. Further, in order to distinguish seed particles from growth particles, the authors report using commercially available diamond grit (0–0.5 m in size). The substrates were then thoroughly cleaned ultrasonically in acetone, rinsed with deionized water, and characterized before growth using SEM imaging and Raman spectroscopy. After run durations of 4 hr, irregularly shaped but faceted particles of size 5–10 m of the type shown in Fig. 25 were observed. The morphology of these particles does not resemble that of diamond crystallites usually observed in other diamond processes (for example, see Fig. 28), suggesting that pure growth is involved here as opposed to the growth and etch processes simultaneously occurring in other CVD processes. Microprobe Raman spectra of these particles (see Fig. 26 showing a representative spectrum) exhibit a beautifully sharp diamond line at 1332 cm⫺1, with no discernible presence of graphitic or amorphous carbon. This work represents the first instance in which diamond growth is observed on unheated substrates by any known process, despite the slow growth rate. In some of their laser photolysis experiments, the H2 stream was dissociated using a tungsten filament resistively heated to about 2000⬚C [68–70]. As expected, the coverage of the particles is substantially increased, as shown in Fig. 27. The morphology of diamond particles produced by photolysis of CO in a hydrogen bath is significantly different from that of crystals normally grown by other CVD techniques. They are irregular in shape and resemble the 0- to 0.5-m diamond grit used for pretreatment of the substrates but are larger (5–10 m in size). The irregularity in shape suggests the following. Because the resolution of SEM imaging is about 50 nm, it is likely that seed particles smaller than 50–100 nm went undetected. These small particles then would have acted as seeds around which diamond could grow, thus explaining the increase in size observed after growth. The typical cubooctahedral shapes (see Fig. 28) and other morphologies usually seen in other
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Schematic of LCVD experiment. (From Ref. 70.)
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Figure 25 surface.
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SEM images of diamond particles grown by LCVD on unheated
Figure 26 Typical Raman spectrum of diamond particles synthesized by LCVD on unheated surfaces.
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Figure 27 SEM image of increased coverage obtained by LCVD, in the presence of a predissociated hydrogen stream.
Figure 28 Typical cubo-octahedral morphology of diamond particles synthesized by hot-filament CVD.
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CVD systems are a result of growth-etch cycles. In conventional CVD systems, the presence of (relatively) large amounts of atomic hydrogen accelerates diamond growth by passivating the growing surface while removing nondiamond carbon and partially etching diamond. An excellent illustration is provided by the interrupted temporal history of growth of diamond around a 100-m seed particle from a CO/H2 gas mixture in a microwave discharge reported in Ref. 72 and reproduced here in Fig. 29. As can be seen, the early stages of growth on the seed crystal give the appearance of an irregular shape for the particle. The morphologies usually observed in other conventional CVD systems begin to become apparent in this case of a CO/H2 plasma only after a long period of time (50 hr) has elapsed [72]. The irregular morphology observed in Refs. 68–70 can then be understood because their experiments spanned only 4 hr. The laser photolysis experiments of Refs. 68–70 provide an excellent example of the use of lasers to drive reactions selectively, specifically those
Figure 29 An excellent illustration is provided by the interrupted temporal history of growth of diamond around a 100-m seed particle from a CO/H2 gas mixture at 45 torr in a microwave discharge reported in Ref. 72. (a) The seed crystal, (b) after 10 hr, (c) after 20 hr, and (d) after 50 hr of growth.
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that result in diamond formation while formation of the graphitic phase is suppressed. The choice of the ArF excimer laser as a source is dictated by the well-known multiphoton process that dissociates CO into C (31P) and ground state O(3P) atoms [73]. Once atomic carbon and oxygen are formed, they react as radicals with the surrounding molecular hydrogen resulting in the formation of desirable CHx species. It is important to note that if H2 were not present, copious amounts of C2, soot, and other nondiamond carbon would result [70]. This channel of C2 production from CO will be addressed further in Sec. V. More recently, it has been shown that optical pumping of CO/Ar gas mixtures with trace amounts of Fe or Ni catalyst results in growth of single-walled carbon nanotubes (see E. Ploenges et al. Carbon Nanotube Production in CO Laser Pumped Carbon Monoxide Plasmas. Paper AIAA2001-0651, presented at the 39th AIAA Aerospace Science Meeting and Exhibit, Reno, Nevada, January 8–11, 2001). In order to understand the laser-driven chemistry in the experiments reported in Refs. 65–67 better, we examine a simple chemical kinetics model of the gas-phase reactions expected to occur in the irradiated CO/H2 mixture. To enable a simple solution, the following assumptions are made. First, every CO molecule in the focal volume is assumed to have dissociated. Second, diffusion and convection are neglected in order to simplify the solution. From the flow rate of CO of about 1 sccm reported in Ref. 68, total pressure of 8 torr, and assuming a local gas temperature of 350 K, the initial number density of CO can be calculated to be ⬃2.2 ⫻ 1021 m⫺3. Similarly, the initial number density of molecular hydrogen based on its flow rate of about 99 sccm is 2.2 ⫻ 1023 m⫺3. Knowing the energy per photon,17 1.0 ⫻ 10⫺18 J and taking the energy per pulse as 50 mJ, the total number of photons in the laser-irradiated volume can be calculated. Assuming this volume to be about 1 mm3, the number density of photons is then approximately 4.8 ⫻ 1025 m⫺3. Because this is much bigger than the initial CO concentration of 2.2 ⫻ 1021 m⫺3 and assuming no more than three photons are required to dissociate a CO molecule, the first assumption is justified. Thus, the initial gas composition within the laser irradiated region can be taken to be nCO = 0 m⫺3, nH2 = 2.2 ⫻ 1023 m⫺3, and nC = nO = 2.2 ⫻ 1021 m⫺3. However, diffusion and flow will serve to replenish the laser-irradiated region with CO and H2 reactants. A more representative and realistic estimate of the gas composition can be obtained by scaling the initial concentrations of atomic carbon and oxygen with respect to the ratio of the irradiated volume to the
17
Because an ArF excimer laser was used in these experiments, the energy of a photon of wavelength 193 nm is (6.626 ⫻ 10⫺34 J-sec) ⫻ (3 ⫻ 108 m/sec)/(193 ⫻ 10⫺9 m) or 1.03 ⫻ 10⫺18 J.
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gas volume adjacent to the substrate. Taking the volume adjacent to the substrate to be about 1 cm3 and the irradiated volume as about 1 mm3, the initial composition for each successive interval is: nCO = 2.2 ⫻ 1021 m⫺3, nH2 = 2.2 ⫻ 1023 m⫺3, and nC = nO = 2.2 ⫻ 1018 m⫺3. With these initial values, the set of rate equations of the form dni = f (ni , nj ) dt
for all j
(12)
describe the evolution of concentration of species i, where ni and nj represent the number densities of the species listed in the reactions given in Table 4.
Table 4 Elementary Reactions and Rate Constants Evaluated at 300 K, Used to Simulate Gaseous Environment Existing Within the Laser-Irradiated Volume Adjacent to the Substrate in the Experiments Reported in Refs. 68–70 Reaction C ⫹ H2 → CH2 C ⫹ H2 → CH ⫹ H H2 ⫹ O → OH ⫹ H C ⫹ CO → C2O CH2 ⫹ H2 → CH3 ⫹ H CH2 ⫹ CO → C2H2 ⫹ O CH2 ⫹ CO → CH2CO H2 ⫹ C2O → CH2 ⫹ CO CH ⫹ H2 → CH3 CH ⫹ H2 → CH2 ⫹ H C2H ⫹ O → CH ⫹ CO CH ⫹ CO → HCO ⫹ C H ⫹ CO → HCO H ⫹ CO ⫹ H2 → CHO ⫹ H2 OH ⫹ CO → H ⫹ CO2 CH2 ⫹ CH2 → C2H2 ⫹ H2 CH3 ⫹ CH3 → C2H6 CH2 ⫹ CH3 → C2H4 ⫹ H CH2 ⫹ CH2 → CH3 ⫹ CH CH ⫹ H → C ⫹ H2 CH3 ⫹ H → CH2 ⫹ H2 C2H2 ⫹ O → CH2 ⫹ CO CO2 ⫹ O → CO ⫹ O2 CO2 ⫹ H → CO ⫹ OH CH3 ⫹ H → CH4
Rate constants 7.11 5.0 7.05 6.31 3.2 2.14 1.0 7.0 2.66 2.66 3.0 2.09 6.92 1.15 1.32 4.98 4.06 6.64 2.30 4.98 3.2 2.14 1.39 1.87 3.32
⫻ ⫻ ⫻ ⫻ ⫻ ⫻ ⫻ ⫻ ⫻ ⫻ ⫻ ⫻ ⫻ ⫻ ⫻ ⫻ ⫻ ⫻ ⫻ ⫻ ⫻ ⫻ ⫻ ⫻ ⫻
10⫺44 m3/sec 10⫺151 m3/sec 10⫺24 m3/sec 10⫺44 m3/sec 10⫺24 m3/sec 10⫺63 m3/sec 10⫺21 m3/sec 10⫺19 m3/sec 10⫺17 m3/sec 10⫺17 m3/sec 10⫺17 m3/sec 10⫺17 m3/sec 10⫺17 m3/sec 10⫺46 m6/sec 10⫺19 m3/sec 10⫺17 m3/sec 10⫺17 m3/sec 10⫺17 m3/sec 10⫺23 m3/sec 10⫺17 m3/sec 10⫺27 m3/sec 10⫺19 m3/sec 10⫺55 m3/sec 10⫺35 m3/sec 10⫺16 m3/sec
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The reactions listed in Table 4 are not meant to be exclusive but represent only one set of reactions thought to be important for the conditions reported in Refs. 68–70. Equations (12) then represent a set of coupled nonlinear equations for ni that can be solved numerically using a method such as that described in Ref. 75. In these calculations, the ArF excimer laser pulses are assumed to remain off for 50 msec (corresponding to a 20-Hz repetition rate) and have pulse durations of 20 nsec. Each pulse interval is hereafter defined as the pulse duration of 20 nsec plus the laser-off time of 50 msec. Compositions of all species (except CO, H2, C, and O) are updated for each interval from values calculated at the end of the previous interval. The compositions after 2 sec or 40 such intervals thus calculated are shown in Figs. 30–32. It can be seen that the carbonaceous species present in order of abundance are HCO, C, CO, CH3, CH2CO, CH4, CH2, C2H6, C2H4, C2H2, and CH. Among these, the least credible is the unusually high amount of atomic carbon present. This must be regarded as a fictitious result because
Figure 30 Compositions of various species calculated using the reactions in Table 4 are shown here after 2 sec or 40 intervals.
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Figure 31 Compositions of various species calculated using the reactions in Table 4 are shown here after 2 sec or 40 intervals.
recombination of C to form either C2 or CO was not included in our reaction set listed in Table 4. Further, the absence of nondiamond carbon either on the SEM images or in the Raman spectra is conspicuous in the experiments reported in Refs. 68–70. This strongly suggests that C2 species must not have been present in abundance in the gas phase because these are known to produce graphitic carbon. Species capable of removing nondiamond carbon or atomic hydrogen or atomic oxygen from CHxOy complexes bound to the growing diamond surface are also present. In order of abundance, they are HCO, C, O, OH, H, CH3, CH2CO, CH2, and CH. Of these, only the high concentration of C should be ignored for the reasons already mentioned. Given the state of present understanding regarding mechanisms of diamond growth [64], we can attempt to make some assessment regarding the identities of possible precursors important in diamond growth by photolysis of CO.
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Figure 32 Compositions of various species calculated using the reactions in Table 4 are shown here after 2 sec or 40 intervals.
Suppose the linear growth rate (i.e., growth rate normal to the substrate surface) is gr. Then, the number of carbon atoms added per unit time per unit area of the growing surface by the carbonaceous precursor, n˙ ⬙C , is given by n˙ ⬙C =
grd mC
(13)
where gr is expressed in units of m/sec, d is the mass density of diamond (3515 kg/m3), and mC is the atomic mass of carbon in kg (12 ⫻ 1.6605 ⫻ 10⫺27 kg). At the very least, for growth to occur, the flux of the carbonaceous precursor must exceed this quantity n˙ ⬙C for growth to continue. Thus, we have 1 ¯ ⱖ n⬙ ˙C nXC 4
(14)
where nX is the gas phase number density of the carbonaceous precursor,
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¯ = (8kBT/mX)1/2 is the random thermal speed. In the expression for and C ¯ kB is Boltzmann’s constant (1.3806 ⫻ 10⫺23 J/K), T is the gas temperature C, in the neighborhood of the growing surface, and mX is the molecular mass of the carbonaceous precursor. In using the expression for the random thermal flux in Eq. (14), it has been assumed that the velocity distribution function of species X is described by the Maxwell-Boltzmann distribution. Using Eqs. (13) and (14), the order of magnitude of the minimum concentration nX necessary to explain an observed growth of gr can be determined. For the specific conditions of the CO photolysis experiments reported in Refs. 68–70, let us take the growth rate to be 1 m/hr, gas temperature T = 300 K, and mX to be on the order of 2 ⫻ 10⫺26 kg. For these values, Eq. (14) yields nX ⬇ 1018 m⫺3. Expected values of the gas composition for times on the order of an hour can be determined from calculated results such as those shown in Figs. 30–32 and compared with this minimum value of nX. Such an order of magnitude estimate shows that candidate growth species in diamond growth by photolysis of CO (excluding C and CH4) are HCO, CO, CH3, and CH2CO. Of these, CO can be eliminated as a likely growth precursor because it is an extremely tightly bound molecule and difficult to dissociate. CH3 emerges as a strong contender as growth species based upon what is known regarding growth mechanisms in other CVD systems [64]. However, the abundant presence of HCO and CH2CO cannot be overlooked, and these may well play a role in diamond growth by photolysis of CO. The other somewhat remote but possible candidate growth species is CH2, as its concentration is lower than that of CH3 only by about two orders of magnitude. Other species such as C2H6, C2H4, C2H2, and CH can be safely eliminated because they are present in substantially smaller amounts. Although the foregoing analysis yields some insight into the identity of the key carbonaceous precursor for diamond growth, it must be emphasized that it is inexact and can be viewed only as an order of magnitude estimate. Nevertheless, such analysis is useful for inferring trends and may suggest process improvements that lead to an increase in yield or growth rate. In this section, we have examined existing photolytic techniques for diamond synthesis. Of significance is the fact that growth of diamond particles has been realized under conditions where the substrate is not heated at all (see Refs. 68–70 on photolysis of CO), and growth of diamond films has been reported at temperatures as low as 450⬚C (see Ref. 63 on photolysis of CCl4). Given the strong interest in lowering growth temperatures [76], these represent significant advancements in the state of the art. It is also of significance to note that in the photolysis experiments involving CO there is evidence that selectivity has been achieved because nondiamond carbon was not codeposited. The absence of nondiamond carbon is conspicuous, and the microprobe Raman spectra show a particularly sharp diamond line
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at 1332 cm⫺1. All these photolysis techniques, however, appear to require predissociation of hydrogen in order to increase growth rates. Of these, the photolytic decomposition of CCl4 in a predissociated hydrogen stream appears closest to commercial realization. An estimate of cost can be made based upon the results reported in Ref. 63. Using their reported growth rate of 1 m/hr but over an area of 1 cm2 (the actual deposition area was not disclosed in Ref. 63) and an input average power of 10 W, the energy cost is approximately 2 ⫻ 107 J/carat for a yield of 4.9 ⫻ 10⫺7 carats/sec.
V.
OPTICAL PUMPING
A.
Introduction
Optical pumping refers to the excitation of the internal modes of atoms or molecules. Light, which is not necessarily coherent, may be used to excite rotational, vibrational, or electronic states of atoms and molecules by what is known as resonant absorption. In contrast to the relatively higher power, multiphoton processes introduced in the earlier sections of this chapter, here we focus on single-photon resonant absorption. The term resonant absorption refers to the fact that there is an excellent coincidence between the frequency of the exciting radiation or light and the energy gap (divided by Planck’s constant) associated with a particular rotational, vibrational, or electronic transition. Further, in contrast to multiphoton processes, optical pumping or resonant single-photon processes are characterized by much lower intensities. Coherence of the exciting radiation is therefore not strictly necessary, although they are needed in practice in order to exceed threshold intensities for absorption. Such threshold absorption levels exist because there are losses due to quenching by collisional or radiative processes that must be overcome. Internal modes of diatomic molecules were introduced briefly in Sec. I. As discussed earlier, these comprise rotational, vibrational, and electronic means of energy storage for molecules. At equilibrium, the average energy in these modes can be described by a single temperature, T. Further, the equipartition theorem18 from statistical thermodynamics tells us that each squared term in the Hamiltonian (corresponding to the degrees of freedom associated with each mode of molecular motion) contributes on average kBT/ 18
The equipartition theorem states that each squared term in the Hamiltonian contributes kT/2 to the average energy. Examples of squared terms are kinetic energy of a particle, which scales as velocity squared and potential energy in a linear spring, which scales as displacement squared. In Hamilton’s formulation of classical mechanics, velocity and displacement are examples of canonical independent variables.
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2 to the average energy of a molecule. For example, there are three degrees of freedom in translation for an atom or molecule. Each appears as a squared term in the kinetic energy arising from the sum of the squares of the three components of velocity (i.e., 12 mC 2x ⫹ 12 mC 2y ⫹ 12 mC 2z ). Hence, the contribution from the translational motion to the total energy of the particle is (3/2)kT. Similarly, for a diatomic particle, the contribution from the rotational modes is kT (corresponding to two degrees of freedom in rotation), and the contribution from the vibrational mode is kT (corresponding to the potential energy and kinetic energy associated with the single degree of freedom in vibration). The essential point here is that the energy in each of these modes is dependent on a single quantity, the gas temperature. When the gas is out of thermodynamic equilibrium, the energy in these modes can be different and is no longer identified by a single temperature. This is of relevance to optical pumping because the resulting situation is one where the energy in the vibrational and electronic modes cannot be described solely by the gas temperature, which characterizes the rotational and translational modes. Equilibration of the energies in the various modes of molecules occurs by collisions. As explained in Sec. I, rotation and translation equilibrate within a few collisions. However, for diatomic molecules such as CO, NO, and N2, it takes many collisions before the vibrational modes equilibrate with the external modes (i.e., rotation and translation). The rate at which energy is exchanged by the various modes via collisions can be described by a rate coefficient, analogously to chemical processes. For example, one can speak of a rate coefficient for vibration-vibration (VV) energy transfer, vibration-translation (VT) energy transfer, vibration-rotation (VR) energy transfer, vibration-electronic (VE) energy transfer, and so on. Thus, each species under such nonequilibrium conditions should be thought of as a specific molecule in a particular vibrational state of an electronic manifold. An example is provided by the energy level diagram of CO shown in Fig. 33, which shows the ground electronic state of CO (X1⌺g⫹) along with its vibrational levels (denoted as the many horizontal lines within the curve). Also shown in Fig. 33 are other electronic states along with their vibrational levels. The rotational levels are not shown for clarity, but they would appear as additional horizontal lines in between successive vibrational levels. For CO at 300 K, the characteristic time for VT transfer19 is expressed as ⬃10 sec-atm for the v = 1 level, while the characteristic time for VV energy transfer is ⬃0.1 msec-torr. At 1 atm, this yields a characteristic time of 10 19
This VT rate is for CO-CO collisions. The corresponding values for CO-He and CO-Ar VT relaxing collisions depleting v = 1 and populating v = 0 are 4 ⫻ 10⫺3 sec-atm and 103 secatm, respectively.
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Figure 33
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Energy level diagram for CO.
sec for VT relaxation for the v = 1 level while its characteristic time for VV exchange is 76 msec. This means that at any given pressure in the collisiondominated regime, vibrationally excited CO will tend to retain energy in vibration for relatively long durations if VT processes are ineffective. Chemical reactions proceed from vibrationally excited or electronically excited states because these promote the necessary breaking of chemical bonds. Optical pumping provides a means of producing vibrationally or electronically excited molecules. This technique has been pioneered by Rich and coworkers [2,77–79], Brechignac et al. [80], and Urban and coworkers [81,82] for studying various nonequilibrium molecular energy transfer processes with applications to electrical discharges, flames, and supersonic expansions. In this section, we discuss in detail the use of this novel technique for diamond synthesis. Although other laser sources and precursor molecules can be targeted using the technique presented in this section, we focus specifically on the technique known as laser-excited CVD or LECVD already reported in the
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literature [70,83,84]. In this method, a CO laser operating on several lines (i.e., transitions) from v = 12 → v = 11 through v = 1 → v = 0, is directed on CO gas molecules in a slowly flowing 95% CO/5% CH4 gas mixture at room temperature. The laser energy is absorbed by the target CO molecules initially in the ground electronic state (X1⌺⫹) and redistributed via anharmonic VV exchange collisions primarily to the internal modes of CH4 as well as to higher vibrational levels within CO. Diamond particles similar in morphology to those produced in Refs. 68–70 were observed on monocrystalline silicon wafers over substrate temperatures ranging from ⬃600⬚C down to room temperature. In order to understand this technique better, it is necessary to provide some detailed background on the optical pumping of CO, a topic on which there exists a considerable amount of literature. At room temperature, most CO molecules are in their ground electronic manifold (X1⌺⫹) with a majority of them occupying the ground vibrational level v = 0. The equilibrium distribution of populations versus vibrational level is described by the well-known Boltzmann distribution: nv = nv=0e⫺(Ev⫺Ev=0)/kT
(15)
where nv is the number density of CO molecules in vibrational level v with energy Ev = ប(v ⫹ 12) ⫺ បe(v ⫹ 12)2, nv=0 is the number density of CO molecules in vibrational level v = 0 with energy Ev=0, and T is the translational mode temperature. In the expression for the vibrational energy Ev, ប = h/2, is the frequency of the molecule as an oscillator, v is the vibrational quantum number, and e is the anharmonicity. Equation (15) states that at equilibrium at a temperature T, most molecules can be found in level v = 0, with the population in higher vibrational levels decreasing exponentially with v. Therefore, if radiation from a CO laser operating on vibrational transitions other than v = 1 → v = 0 is directed on gaseous CO at room temperature, no absorption will result and the gas will be transparent to the laser irradiation. It is evident that for cold CO gas to absorb the radiation, either the laser must have some output on the lowest transition v = 1 → v = 0 (i.e., shortest wavelength) or the gas must be heated (so that upper vibrational levels have sufficient population to absorb the longer wavelength laser transitions). Commercially available CO lasers typically do not have output on this shortest wavelength transition, and therefore such a CO laser must be specially built [70,83]. When laser output on the v = 1 → v = 0 transition is present, CO gas molecules can now absorb this laser energy and populate the v = 1 level. These molecules in level v = 1 now absorb the laser output from the v = 2 → v = 1 transition and populate the v = 2 level, and so on. In this manner, CO molecules can be optically pumped up to vibrational level v = 12, or whatever may be the vibrational level corre-
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sponding to the longest wavelength laser transition. Collision-induced vibrational energy (VV) exchange processes then proceed according to CO(v) ⫹ CO(w) S CO(v ⫺ 1) ⫹ CO(w ⫹ 1),
wⱖv
(16)
This collisional vibrational energy exchange is favored due to the anharmonicity of molecules20 and serves to populate vibrational levels much higher than those initially populated by the laser irradiation. Levels as high as v = 40 in the ground electronic manifold of CO (corresponding to energies of about 10 eV) have been populated in this manner. This laser excitation followed by redistribution of the absorbed laser energy among the higher vibrational levels is known as anharmonic VV pumping. The resulting population distribution, henceforth referred to as the TRR distribution, is highly non-Boltzmann and was first derived by Treanor, Rich, and Rehm in 1968 [85]. For the limiting case of VV exchange collisions dominating VT exchange, the TRR distribution is [85]: nv = nv=0 exp
再
冎
(ប v ⫺ 2បev) បev(v ⫺ 1) ⫺ kT kT1
(17)
where T1 is the ‘‘vibrational temperature’’ of level v = 1 defined by nv=1 = nv=0e⫺ប(1⫺2e)/kT1 or ប(1 ⫺ 2e) T1 = ⫺ nv=1 k ln nv=0
冉 冊
(18)
The form of Eq. (17) for the TRR distribution can also be derived from statistical thermodynamics in a manner analogous to that used to derive the Bose-Einstein or Fermi-Dirac distribution. This is done by maximizing the thermodynamic probability subject to the constraints that the total number of particles, total energy, and total number of vibrational quanta per molecule are fixed (i.e., conserved) [85]. Figure 34 shows a semilog plot of the normalized number density or population density as a function of vibrational quantum number for several cases. These are (1) the Boltzmann distribution as given by Eq. (15) but neglecting anharmonicity in the molecule, (2) the Boltzmann distribution 20
Anharmonicity refers to the fact that the vibrating bond between C and O atoms in the CO molecule exhibits departure from that of a harmonic oscillator. As a consequence, the gap between successive vibrational energy levels decreases as the vibrational quantum number (or vibrational energy) increases.
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Figure 34 Schematic showing a log-linear plot of normalized populations of molecules in vibrational level v versus vibrational quantum number (or vibrational energy) for three cases: (1) equilibrium, where the variation is given by the Boltzmann distribution [Eq. (15)] but without the effects of anharmonicity; (2) equilibrium, where the variation is given by the Boltzmann distribution [Eq. (15)] but including the effects of anharmonicity; and (3) nonequilibrium in which the average vibrational energy is different from the average translational energy.
including anharmonicity, and (3) the limiting TRR distribution as given by Eq. (17), with Tv ≠ T. It is important to point out that the concept of a vibrational temperature Tv that represents a meaningful average energy in vibration makes sense only when the variation of normalized number density versus vibrational quantum number is log linear. When departure from log linearity occurs, a single vibrational temperature cannot be defined for all vibrational levels. This is what happens in the case of the Boltzmann distribution including the effects of anharmonicity with Tv ≠ T or in the case of the ideal TRR distribution. The vibrational quantum number at which the number density attains a minimum for the TRR distribution can be determined by differentiating ln(nv /nv=0) from Eq. (17) with respect to v and setting it equal to zero. The result yields what is known as the Treanor minimum: vmin =
(1 ⫺ 2e) 2e
冉冊 T T1
⫹
1 2
(19)
A similar minimum exists for the Boltzmann distribution (i.e., equilibrium distribution) when the effect of molecular anharmonicity is included: (vmin)equilibrium =
(1 ⫺ e) 2e
(20)
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Values of frequencies and anharmonicities for some diatomic species are given in Table 5. As can be seen, the anharmonicity, e, takes on values that are usually less than or equal to 0.01. This means that (vmin)equilibrium ⬇ 1/2e ⬇ 50. Such high vibrational levels are so close to the dissociation limit that the quadratic behavior of anharmonicity is never observed under equilibrium conditions. Thus, at equilibrium, the populations of vibrational levels are nearly log linear with vibrational quantum number below the dissociation limit, even when the effect of anharmonicity is included. In contrast, the TRR distribution yields vmin ⬇ (82)(1/5) ⬇ 16 for CO for T1 /T ⬇ 5. Thus, for the species listed in Table 5, the TRR distribution yields a minimum in the number density variation when the vibrational quantum number is on the order of 10 and the vibrational temperature of level v = 1 is a few times higher than the translational or gas temperature. In reality, the TRR distribution is modified by the competing effect of VT relaxation. VT energy transfer scales with vibrational quantum number and hence becomes important at the higher quantum numbers. The degree to which vibrational nonequilibrium persists in a gas therefore depends on the competing strengths of VV and VT or other loss processes such as spontaneous emission. In molecules such as CO and NO, spontaneous emission is slower than VT processes at pressures of interest. In homopolar diatomic molecules such as O2 and N2, there is no spontaneous emission because these molecules possess no natural or intrinsic dipole moment. We therefore illustrate the realistic VV up-pumping process by considering VV and VT processes alone, with the rotational modes in equilibrium with the translational modes at temperature T. Consider a diatomic molecule A2 undergoing the following single-quantum processes: VV energy transfer: k w⫺1,w v⫹1,v A2(v ⫹ 1) ⫹ A2(w ⫺ 1) → A2(v) ⫹ A2(w) ← k w,w⫺1 v,v⫹1
Table 5
(21)
Frequencies and Anharmonicities for Some Diatomic Molecules
Parameter ប (cm⫺1) បe (cm⫺1) ប/2បe Source: Ref. 86.
CO
H2
N2
NO
O2
OH
2169.8 13.3 82
4401.2 121.3 18
2358 14.1 84
1904 14 68
1580 12 66
3735 83 22
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VT energy transfer: kv⫹1,v A2(v ⫹ 1) ⫹ M → ← A2(v) ⫹ M kv,v⫹1
(22)
where the VV and VT rates are given in Refs. 75 and 79. The evolution of the population of level v is then given by the following rate equations, including spontaneous emission as a loss mechanism [2]: dnv = VTv ⫹ VVv ⫹ SRDv dt
(23)
where the terms on the right-hand side represent net VV transfer, VT transfer, and spontaneous radiative decay into level v. Expressions for each of these terms, for CO, are given in Ref. 75. Process (21) is shown schematically in Fig. 35 for two general molecules A and B. Now, if E(v) is the energy of molecules in vibrational level v, then E(v ⫹ 1) ⫺ E(v) > E(w) ⫺ E(w ⫺ 1) for w ⱖ v because of the effect of anharmonicity. Thus, if we treat the right-hand side of process (21) as the ‘‘products’’ and the left-hand side as the ‘‘reactants,’’ we have ⌬E = E(v) ⫹ E(w) ⫺ E(v ⫹ 1) ⫺ E(w ⫺ 1) = [E(w) ⫺ E(w ⫺ 1)] ⫺ [E(v ⫹ 1) ⫺ E(v)] so that ⌬E ⱕ 0 for w ⱖ v. ⌬E is called the resonance defect. This shows that the forward process of (21) is favored at lower temperatures because it is exothermic in the sense that the energy difference ⌬E flows out of the vibrational modes into the external modes of rotation and translation for w ⱖ v. On the other hand, the reverse of process (21) is endothermic in the sense that energy must be provided from the external modes to sustain it. This can be expected to occur only at high temperatures. Energy can be retained in the vibrational modes, therefore, if an external source (such as a laser or optical pump) provides energy to the vibrational mode
Figure 35 Schematic showing intermolecular vibrational energy transfer, preferentially transferring vibrational energy from the molecule with the more widely spaced energy levels to the molecule with the more closely spaced energy levels [2]. When both collision partners are molecules of the same species, then such VV transfer is called intramolecular vibrational energy transfer. The difference between ⌬v and ⌬w is called the resonance defect.
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to overcome loss processes such as process (22) or else process (22) is inhibited or minimized. Process (22) scales with vibrational quantum number v and temperature; i.e., the rate for vibrational deexcitation by VT transfer increases with increasing v and increasing T. Hence, as long as the energy corresponding to the resonance defect is removed either by having the gas flow or by having a suitable diluent such as argon, a low gas temperature can be maintained or VT losses can be lowered compared with VV exchange. Under these conditions, the solution of the rate equations (23) yields a fully VV pumped distribution and is indeed observed as shown in Fig. 36. The fully VV pumped distribution is observed in optically pumped gases as well as in common plasmas. It is, in fact, ubiquitous in plasmas containing tightly bound molecular species such as CO, NO, N2, or O2. In such plasmas, the mean electron energy is lowered because inelastic collisions between electrons and molecules dominate all other processes. Consequently, the electron (translational) energies in such discharges strongly resemble the fully VV pumped distribution shown in Fig. 36. The result is that a single electron temperature cannot be identified, and hence we must speak of electron energies and a corresponding electron energy distribution function or EEDF that is highly non-Boltzmann. These electron impact processes can be included in the rate equation formulation as given by Eq. (23) and are crucial in determining the chemical composition in a plasma. Process (16) is an example of intramolecular energy transfer, or transfer of vibrational energy between molecules of the same chemical species. Intermolecular energy transfer, or transfer of vibrational energy between molecules of different chemical species, also occurs. In both cases, energy is transferred preferentially from the molecules with widely spaced vibrational energy levels to those with more closely spaced energy levels. LECVD takes advantage of this point precisely. Further, this energy transfer is a result of molecular collisions and scales with pressure (or density) as long as relatively low external mode (i.e., translational and rotational) temperatures are maintained or as long as rates of population of vibrational levels by VV exchange collisions are favorable compared with the rates of depopulation of these levels by VT collisional processes. The effect of anharmonicity in a diatomic or small polyatomic molecule is such that any vibrational energy over and above the equilibrium value (corresponding to the translational temperature T ) is rapidly redistributed to higher vibrational states. Thus, molecular states at high vibrational levels can be populated, especially when the gas temperature is very low. The process, however, is not restricted to low gas temperatures and has been shown to occur at elevated gas temperatures on the order of 1500 K [87].
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Figure 36 Fully VV pumped distribution from Ref. 79 measured from an optically pumped mixture of 1% CO/Ar at a total pressure of 100 torr. Note that at low vibrational quantum numbers, the VV pumped distribution resembles the TRR distribution but instead of displaying a minimum exhibits a plateau known as the TRR plateau characterized by v ln(nv=0/nv) = constant. The distribution exhibits a sharp drop at the higher vibrational quantum numbers because the VT rates overwhelm the VV rates at high values of v.
B.
Optical Pumping of Gases (Laser-Excited CVD or LECVD) for Diamond Growth
Optical pumping of gases using a low-power laser can, under the right conditions, result in production of molecules with high vibrational energies. These molecules can then be made to react among themselves or transfer energy to other species to excite them to specific vibrational or electronic states and hence serve as huge heat baths for driving chemical reactions. A particularly dramatic demonstration of the power and usefulness of this tech-
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nique is the early work of Kildal and Deutsch [88]. At the time, laser sources in the mid-IR such as the CO laser were not as well developed as they are today. Consequently, Kildal and Deutsch [88] went to elaborate lengths to obtain some coherent output at ⬃4.7 m by frequency doubling the 9.4 m laser line from a CO2 laser. The resulting laser output at 4.7 m is coincident with the lowest vibrational transition, v = 0 → v = 1, in CO so that CO can be vibrationally excited to level v = 1 at room temperature. Once this state is populated, of course, other higher vibrational levels can be populated via anharmonic VV pumping: CO(v = 1) ⫹ CO(v = 1) → CO(v = 0) ⫹ CO(v = 2), followed by CO(v = 1) ⫹ CO(v = 2) → CO(v = 0) ⫹ CO(v = 3), and so on. Kildal and Deutsch used this approach to optically pump various gas mixtures containing CO, where CO is used as the energy ‘‘donor’’: CO/ C2H2, CO/C2H2 /H2, CO/CO2, CO/N2O, CO/CS2, CO/CS2 /H2 and CO/OCS [88]. Further, by placing mirrors on the chamber containing these gas mixtures, Kildal and Deutsch were able to obtain lasing from the ‘‘acceptor’’ molecules in which population inversions (partial or total) were sustained by energy supplied from the optically pumped and subsequently vibrationally excited CO molecules. The fact that such transfer lasers could be produced was testament to the rapidity of intermolecular vibrational energy transfer versus loss of this energy to translational modes. A laser source at 4.7 m (i.e., laser output on the CO fundamental v = 1 → v = 0) would have made the optical pumping of CO much easier. Similar energy transfer and VV pumping have been previously reported in CH4, CD3H, CD4, CH2D2, and CH3F [89]. The CO laser was invented as early as 1965 by Patel [90], but theoretical understanding of its operating principle would await the development of the theory of anharmonic VV pumping developed by Treanor, Rich, and Rehm in 1968 [85]. This theoretical development subsequently led to more powerful supersonic and cryogenically cooled CO laser sources [77–79]. Today, the CO laser is the most efficient gas laser, displaying efficiencies near 50%. Steady development in the 1980s and 1990s led to compact cryogenically cooled CO laser sources capable of significant output on the lowest vibrational transition (v = 1 → v = 0). This has been important in enabling the optical pumping of CO at room temperature. We begin with a description of the optical pumping of CO diluted in argon. It is necessary to understand optical pumping in this simple mixture before proceeding with a description of the LECVD process that utilizes the optical pumping of other cold CO-containing gas mixtures. An excellent review can be found in Ref. 2. Consider a mixture of CO/Ar (typically a few percent CO with the balance being argon) flowing at 100 sccm, total pressure of about 50 torr, and temperature of 300 K, irradiated by a CO laser operating multiline with some of its output on the lowest transition (v
Laser-Assisted and Optical Pumping Techniques
387
= 1 → v = 0). Usually, lasing on the lowest transition is difficult but can be relatively easily produced using an intracavity Q switch [70,82]. This usually results in about 100 mW output on the lowest transition, with a total output of about 5 W on all lines. With improved designs [91], cw (continuous wave) output with several watts on the lowest transition can be obtained. The CO in the CO/Ar mixture absorbs this radiation on the v = 1 → v = 0 transition, which populates the v = 1 energy level above its equilibrium value. This in turn enables laser output on the v = 2 → v = 1 transition to be absorbed, populating the v = 2 level above its equilibrium value, and so on. In this manner, levels v = 9 down to v = 1 are populated by absorption of energy from the laser beam. A typical spectrum of the CO laser emission is shown in Fig. 37. It is important to highlight the fact that this absorption process is a resonant, single-photon process and therefore does not require high laser powers or pulse energies. This absorbed laser energy is then redistributed to the higher vibrational levels by anharmonic VV pumping. The result is a fully VV pumped distribution as shown in Fig. 36. Note that only one por-
Figure 37 Typical spectrum showing output from a CO laser. This laser is capable of operating at about 10 W cw while also lasing on the v = 1 → v = 0 transition (leftmost peak). Each successive set(s) of peaks represents higher vibrational bands from v = 2 → v = 1 up to v = 10→ v = 9. (From Ref. 91.)
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tion of the TRR distribution is evident because VT rates scale or increase with vibrational quantum number and prevent the total inversion that is characteristic of the TRR distribution. Instead of a minimum, the real VV pumped distribution exhibits a near-plateau region over which vnv ⬇ constant, followed by a steep drop in the vibrational populations at the higher vibrational quantum numbers. The reason for both trends is that both VV and VT rates increase with v, but eventually VT rates exceed the VV rates that drive the up-pumping at the higher vibrational quantum numbers. A fully VV pumped vibrational population distribution in a flowing CO/Ar mixture results in population of vibrational states as high as v = 42, corresponding to energies of more than 8 eV [79]. This is significant because the dissociation energy of the CO molecule is about 11 eV. When a molecule such as CO is pumped to vibrational levels as high as v = 42, quite a bit of energy is engaged within the vibrational mode. This energy should be available to drive chemical reactions or for transfer to other modes of energy storage such as electronic excitation. Indeed, a host of energy transfer processes occur such as those that result in population of low-lying electronic states by VE (vibration-to-electronic) transfer [79,82]: CO(X1⌺⫹, v) ⫹ CO(X1⌺⫹, w) → CO(v⬘) ⫹ CO(a3⌸, w⬘), Ev ⫹ Ew > 6 eV ⫹
⫹
⫹
CO(X ⌺ , v) ⫹ CO(X ⌺ , w) → CO(X ⌺ , v ⬘) ⫹ CO(A ⌸, w⬘), Ev ⫹ Ew > 8 eV 1
⫹
1
1
CO(X ⌺ , v ⬇ 27) ⫹ M → CO(a ⌸, w⬘) ⫹ M) 1
⫹
3
CO(X ⌺ , v ⬇ 40) ⫹ M → CO(A ⌸, w⬘) ⫹ M 1
1
(24)
1
(25) (26) (27)
Processes (26) and (27) represent direct collisional transfer to electronic states, one of which is the a3⌸ state, which is metastable with a radiative lifetime of 4.4–9.5 msec. Another electronic state, A1⌸, has a radiative lifetime of 10 nsec and therefore serves to limit the vibrational up-pump. It must be pointed out that around v = 42, 2[E(v = 42) ⫺ E(v = 40)] ⬇ [E(v = 1) ⫺ E(v = 0)], so that the following near-resonant process: CO(X1⌺⫹, v ⬇ 42) ⫹ CO(X1⌺⫹, v = 0) → CO(X1⌺⫹, v ⬇ 40) ⫹ CO(X1⌺⫹, v = 1) can occur with high rapidity, as does VT transfer (which scales with the vibrational quantum number v). While this is possible, the VE transfer via process (27) has been observed experimentally, and radiative emission from CO(A1⌸, w) → CO(X1⌺⫹, v) has been observed only after CO(X1⌺⫹, v ⬇ 42) is populated [79]. Thus, despite attaining energies near the dissociation limit, CO is not observed to dissociate because of the rapid VE transfer to
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389
the A1⌸ electronic state, followed by fast radiative decay to the ground electronic state or manifold [79]. Thus, infrared radiation goes into exciting the gas, while ultraviolet light comes out! In contrast, the a3⌸ by virtue of its relatively long lifetime suffers collisions and is either deexcited (collisionally quenched) or reacted.21 Processes (24) and (25) represent energy transfer by an associative mechanism referred to as vibrational pooling. Such processes have high probability of occurrence because they typically involve many combinations of molecular collision pairs with energies that add up to the required energy of a given excited electronic state. Such associative processes have been shown to be instrumental in producing ionization by vibrational pooling and may constitute a major source of ionization in CO-containing electrical discharges [92]: CO(X1⌺⫹, v) ⫹ CO(X1⌺⫹, w) →
再
CO⫹ ⫹ CO(X1⌺⫹, w⬘) ⫹ e⫺, ⫺ (CO)⫹ 2 ⫹ e
Ev ⫹ Ew ⱖ 13.6 eV
(28)
The lower pathway of process (28) resulting in the formation of the dimer ion is favored because it has a slightly lower ionization potential than CO. Regardless, the production of CO⫹ or (CO)⫹ 2 dimer ions usually leads to cluster ions of the form (CO⫹ ⭈ CO⭈ Cn) where n = (0, . . . , 11) [93]. In addition to the aforementioned energy transfer and associative ionization processes, chemical reactions such as the Boudouard reaction occur: {CO(X1⌺⫹, v) or CO(a3⌸, v⬘)} ⫹ {CO(X1⌺⫹, w) or CO(a3⌸, w⬘)} → C* ⫹ CO2, Ev ⫹ Ew > 6 eV
(29)
followed by: C ⫹ CO ⫹ M → C2O
(30)
C ⫹ C ⫹ M → C2 ⫹ M
(31)
The formation of C2 in the gas phase is generally thought to be undesirable for diamond formation. There is some work however, that suggests C2 could be a growth species for nanocrystalline diamond [94,95]. Past optical pumping experiments involving sufficient vibrational excitation of CO to drive the Boudouard reaction (29) leading to C2 formation via (31) have shown formation of graphitic soot [96]. Reaction (31) is one of the channels for Alternatively, there can be collision-induced E-V transfer from CO(a3⌸) into the CO(X1⌺⫹, v ⬇ 27) state, although it has not yet been observed experimentally by optical pumping.
21
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formation of soot in electrical discharges containing substantial amounts of CO. In a plasma, free carbon is produced by electron impact dissociation of CO as well as by reactions (29) and (31) where the appropriate CO(a3⌸) or vibrationally excited CO(X1⌺⫹) is produced by electron impact processes. In contrast, formation of C2 in the gas phase in an optically pumped mixture first requires the formation of atomic carbon via reaction (29), which has an activation energy of about 6 eV. Therefore, it is clear that formation of CO in the a3⌸ state [whose lowest vibrational level is isoenergetic with the v ⬇ 27 state of CO(X1⌺⫹) with an energy on the order of 6 eV] or CO in vibrational levels beyond v ⬇ 13 in the X1⌺⫹ state must be suppressed. The requirement that CO in vibrational levels beyond v ⬇ 13 in the X1⌺⫹ state not be populated stems from the fact that molecules in these levels can pool their energies to form some molecules in v ⬇ 27 of the X1⌺⫹ state or in the low vibrational levels of the a3⌸ state. In general, light molecules or molecules with many rotational modes are very effective in VT relaxing vibrationally excited CO [2,3]. Molecular hydrogen and water are two examples of strong VT relaxers [2,75]. Hence, it is very difficult to produce reactive vibrationally excited CO molecules by optical pumping in mixtures predominantly comprising H2 or H2O, which is typical of growth mixtures used in diamond synthesis by CVD. Rebello et al. circumvented this problem by using a mixture containing mainly CO but with some CH4, which is commonly used as a carbonaceous precursor in CVD processes for diamond synthesis. Diamond growth was reported from a gaseous mixture of 95% CO and 5% CH4, where the CO was optically excited using a Q-switched CO laser operating with its shortest wavelength output on the v = 3 → v = 2 transition [70,83]. These experiments were the first to demonstrate that diamond growth was possible without the need to hyperdilute the gas phase with hydrogen or some other etchant. Such a mixture has never before been reported to yield diamond growth by any other CVD process (HFCVD, plasma CVD, or flame synthesis), and may therefore signal the existence of alternative routes to diamond synthesis. We review their experiment here in detail. A schematic of the LECVD or laser-excited CVD experiment is shown in Fig. 38. A CO laser Q-switched at 200 Hz with the lasing lines shown in Fig. 39 was used to optically pump a flowing mixture of 95 sccm of CO and 5 sccm of CH4 at a total pressure of 50 torr. Of the 2.3 W output (on all lines) from the Q-switched CO laser, approximately 1 W was reportedly absorbed by the CO in the gas mixture [70,83]. Single-crystal Si(100) wafers approximately 2 cm2 in size were used as substrates and heated to a temperature between 600 and 700⬚C. The substrates were supported on a heated holder within the reactor and oriented upside down [83]. This was done to
Laser-Assisted and Optical Pumping Techniques
Figure 38 Ref. 83.)
391
Schematic of the LECVD or laser-excited CVD experiment. (From
take advantage of the buoyancy-driven or free convective flow induced by substrate heating, which tends to drive the gas upward. Such a means of mounting was also used in experiments where the substrates were not heated [70]. The substrates were physically abraded using diamond particles (4- to 8-m sized particles in the earlier work where the substrate was heated [83] and with particles 0–0.5 m in size in the later work where the substrates
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Figure 39 Spectrum showing the lasing lines of the CO laser used in the LECVD experiments. This laser was Q switched at 200 Hz. (From Ref. 83.)
were not heated [70]) suspended in glycerol and subsequently cleaned in deionized water, acetone, and methanol, respectively. The authors reported that such a pretreatment was used to possibly reduce nucleation times and thereby enable short-duration experiments, as is typically done with other diamond growth techniques [70,83]. The substrates were then thoroughly characterized prior to growth using SEM imaging, and Raman spectroscopy (survey as well as with a microprobe). The substrates contained some residual abrasive particles only in a few isolated cases. However, these proved to be debris of the silicon substrate when analyzed using energy dispersive x-ray spectroscopic (EDS) analysis. Further, these debris particles did not show the characteristic diamond line at 1332 cm⫺1 in the corresponding Raman spectra. The authors carefully characterized the substrates before growth and reported on the formation of additional particles approximately 10 m in size clustered together in groups that covered regions on the order of 50 or 60 m2. Figure 40 shows SEM images of the substrate surface before and after growth. The fact that these growth particles were grouped together enabled Raman spectra to be obtained using a microprobe. A representative Raman spectrum and EDS spectrum are shown in Figs. 41 and 42, respectively. Of
Laser-Assisted and Optical Pumping Techniques
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Figure 40 SEM images before and after growth in the LECVD experiment. (From Refs. 70, 83.)
significance is the fact that these experiments have been repeated under conditions where the substrate is not heated and similar diamond growth is observed [70]. Evidently, diamond growth by LECVD does not require the monocrystalline silicon substrate to be heated. During the LECVD growth process, there was no visible emission from the reactor chamber even when the substrate was heated because the resistively heated tungsten filament was hidden from view. However, there was radiative emission from the gas in the infrared. Figure 43 shows a low-resolution IR spectrum from the gas
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Figure 41 Typical Raman spectrum of the diamond particles synthesized by LECVD. (From Ref. 70.)
phase for the case where the substrate was not heated, displaying the overtone emission (i.e., ⌬v = 2) from vibrationally excited CO and emission from the 3 mode of CH4. The 3 mode is the asymmetric stretch mode where a single hydrogen atom vibrates against the CH3 group in the methane. An energy level diagram for CH4 is shown in Fig. 44 along with the corresponding low-lying CO vibrational energy levels. In addition, faint emission was observed from the 4 or bending mode of CH4 [70]. The 4 mode emission is particularly difficult to detect because it overlaps with ambient blackbody radiative emission at 300 K. These spectra serve to verify that CO is vibrationally excited and transfers energy to CH4 to excite it vibrationally as well. The population distribution in CO(X1⌺⫹) can be inferred by comparing the experimentally measured radiative emission with synthetic spectra and is shown in Fig. 45 [70]. It can be seen that the CO in the CO/ CH4 mixture under conditions of interest to diamond growth exhibits vibrational nonequilibrium but is not fully VV pumped. Using the slope of the population distribution for the lower vibrational levels in Fig. 45, a vibrational temperature T = 2200 K can be inferred.22 There are some interesting similarities and differences between these diamond crystals and those observed in other CVD systems. In the first 22
Note that here the translational temperature T ⬇ 300 K ≠ T = 2200 K.
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Figure 42 Energy-dispersive x-ray spectroscopic (EDS) analysis spectrum of the diamond particles synthesized by LECVD. (From Ref. 70.)
place, the morphology of these particles is clearly irregular but still faceted. It resembles the morphology of the particles obtained by photolysis of CO in the CO/H2 mixtures in LCVD (see Sec. IV and Refs. 68 and 69). There is also a strong resemblance of the morphology of the growth particles to that of the high-pressure diamond grit used in the abrasion pretreatment, except the growth particles are larger. The well-defined crystal habits displayed in diamond synthesized by other CVD techniques are due to the simultaneous presence of growth and etch channels. If pure growth occurred without the presence of any etch mechanisms, morphologies such as those in Figs. 25 and 29 would be expected, as discussed earlier in Sec. IV (see Figs. 25 and 29, associated discussion, and Ref. 72). The Raman spectra reveal that the growth particles are indeed diamond as can be seen from the sharp 1332 cm⫺1 line in Fig. 41. The fact that no graphitic lines (at 1580 cm⫺1 and/or at 1343 cm⫺1) are observed in the Raman spectrum is indeed conspicuous. Codeposition of nondiamond (graphitic) carbon is usually ob-
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Figure 43 Low-resolution IR spectrum from the gas phase, displaying the overtone emission (i.e., ⌬v = 2) from vibrationally excited CO and emission from the 3 mode of CH4. The 3 mode is the asymmetric stretch mode where a single hydrogen atom vibrates against the CH3 group in the methane. (From Ref. 70.)
served in diamond growth using other types of CVD processes and is clearly absent in LECVD. The EDS spectrum shows the presence of a small peak corresponding to oxygen. As this peak is observed only in diamond films grown using other CVD techniques in which oxygen or oxygen-bearing compounds are used in the growth gases [15–17,69], this suggests rather strongly that CO may play a role in diamond growth by LECVD. Such an oxygen peak is never observed in EDS spectra of diamond films exposed to ambient air. Therefore, it appears highly likely that oxygen is present on the diamond surface itself as a direct result of the specific LECVD growth environment and may in fact provide clues to growth mechanisms. Control experiments were also done to determine what the growth species might be [70,83]. Substrates having undergone the same abrasive pretreatment, and either heated to the same temperatures or unheated, were exposed to optically pumped 95% CO/5% H2 and 100% CO gaseous environments but showed no presence of diamond. These control experiments suggest that CH4 is the likely carbonaceous precursor. However, the CO
Laser-Assisted and Optical Pumping Techniques
397
Figure 44 Energy level diagram for CH4, along with the corresponding low-lying CO vibrational energy levels. (From Ref. 70.)
could have played a facilitating role as can be inferred from EDS analysis indicating the presence of oxygen on the diamond surface. Unfortunately, experiments involving 13CH4, which would have been more illuminating, were not done.23 As mentioned earlier, the same optical pumping technique has been applied to conditions where the substrate is not heated [70,84]. Earlier experiments [83] were conducted with a heated substrate and v = 3 → v = 2 23
The use of 13CO in the gas phase would have been problematic because the CO laser source would also have had to operate using the same isotope in order to take advantage of resonance absorption.
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Figure 45 Population distribution in CO (X1⌺⫹) inferred by comparing the radiative emission with synthetic spectra. (From Ref. 70.)
as the shortest wavelength lasing transition from the CO laser. As discussed earlier in Secs. V.A and V.B, there is a negligible population of CO molecules in the v = 2 state at room temperature so that laser absorption to trigger the VV pumping could occur only at an elevated temperature. In the more recent work [70,84], however, the CO laser was operated with some output on the v = 1 → v = 0 transition so that CO molecules at room temperature could be vibrationally excited. Indeed, as in the LCVD or laser photolysis technique discussed in Sec IV, diamond growth was obtained on unheated surfaces [70,84]. It is of significance that two completely different laserbased techniques (LCVD and LECVD) involving two completely different gas mixtures (1% CO in 99% H2 for photolysis and 5% CH4 in 95% CO for LECVD) yielded diamond growth on unheated surfaces. This also suggests that growth mechanisms in these completely different experiments may share common features. We shall address the subject of mechanisms shortly. Optical pumping experiments have also been conducted in mixtures of 5% C2H2 in CO [70,84]. A low-resolution emission spectrum from this experiment is shown in Fig. 46 and displays the 3 band of acetylene and faint emission from possibly the combination bands 2 ⫹ 15, 2 ⫹ 2 05 ⫺ 15, 2 ⫹ 2 14 ⫹ 15 ⫺ 14, 3 ⫺ 14, and 1 ⫺ 15 [97]. The spectrum ascribed to the
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Figure 46 Low-resolution emission spectrum from optical pumping of CO/C2H2 mixture showing the 3 band of acetylene and faint emission from possibly the combinations bands 2 ⫹ 15, 2 ⫹ 2 05 ⫺ 15, 2 ⫹ 2 14 ⫹ 15 ⫺ 14, 3 ⫺ 14, and 1 ⫺ 15 [97]. (From Ref. 70.)
3 band may also contain some emission from the 2 ⫹ 14 ⫹ 15 combination band. The 3 band of acetylene arises from its asymmetric stretch mode wherein a hydrogen atom vibrates against the associated C2H. As in the case of the optically pumped CO/CH4 mixture, the vibrational populations of CO can be inferred from the overtone emission spectrum. This is shown in Fig. 47 for the CO/C2H2 mixture, and appears to indicate that levels up to about v = 9 or v = 10 are populated. However, no diamond growth was obtained from optically pumped CO/C2H2 mixtures despite the obvious energy transfer from CO to acetylene [70]. The LECVD experiments reported in Refs. 70, 83, and 84 were independently reproduced in the Laboratoire d’E´nerge´tique Moleculaire et Macroscopique, Combustion (E.M2.C.) at E´cole Centrale Paris in France [98]. These results confirm diamond growth both on unheated silicon substrates as well as on heated substrates [98]. Additional experiments in which high-resolution spectra were obtained are reported in Ref. 98 as well. Figure
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Figure 47 Vibrational populations of CO inferred from the overtone emission spectrum are shown here for the CO/C2H2 mixture and indicate that levels up to about v = 9 or v = 10 are populated. (From Ref. 70.)
48 shows a high-resolution spectrum of the same bands of CO and CH4 that are displayed in Fig. 43 for the experiments reported in Refs. 70, 83, and 84. With 4 W of the incident 10.8 W cw laser power (on all lines) absorbed, Ref. 98 reported that emission measurements indicate that the CO alone (without addition of CH4) is pumped up to levels near v = 17, which corresponds to energies of about 4 eV. Upon addition of the CH4, emissions from CO molecules are observed only up to about v = 10, which corresponds to energies of about 2.5 eV. Note that in the latter case, formation of C2 in the gas phase is suppressed because reaction (29) requires an activation energy of about 6 eV. Reference 98 also presents detailed analyses using high-resolution emission and absorption spectroscopy. Careful analysis shows that the spectrum largely attributed to the 3 mode of CH4 actually contains other lines arising from the combination bands 3 ⫹ 4 → 4 and 3 ⫹ 2 → 2, also known as the ‘‘hot bands’’ of methane. From these spectra, a vibrational temperature T of 1950 K was inferred [98]. Of significance is the fact that the detailed gas-phase spectroscopic studies of Ref. 98 indicate the presence only of CO and CH4 in both emission and absorption measurements. No evidence was found for any other chemical species or reaction by-products, at least within detection limits. Clearly, either re-
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Figure 48 High-resolution spectrum of the overtone bands of CO and 3 mode emission from CH4 from the E´cole Centrale Paris experiments. (From Ref. 98.)
action by-product concentrations were too small to be observed or reactions occurred on the growth surface. Diamond growth on unheated, seeded substrates has been reported using two completely different laser-based techniques (laser photolysis and LECVD) and using two very different gas mixtures (1% CO in H2 in the photolysis case and 5% CH4 in CO in the optical pumping case). These results beg the question of whether or not the two techniques share some commonality in growth mechanisms. At first glance, it appears they share
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no common features. Laser photolysis uses a ‘‘conventional’’ gas mixture for diamond growth and therefore can be explained using present understanding of diamond growth by CVD processes [64], which identifies the CH3 radical as the growth precursor. Once CH3 groups are added to a growing diamond surface, the atomic hydrogen is removed by the atomic hydrogen present in the gas phase. In contrast, LECVD uses a gas mixture that normally would not yield diamond growth if used in other conventional non–laser-based CVD methods. There cannot be an appreciable amount of atomic hydrogen in the LECVD environment, or else the CO would not VV pump [2,75]. Because the particle growth rates for the two methods are comparable (on the order of 1 m/hr) as are the resulting morphologies and Raman spectra, we expect the carbonaceous precursor to be present in concentrations on the order of 1018 m⫺3 following the same reasoning as in Sec. IV. It is extremely difficult to infer growth mechanisms or detailed channels based upon the experiments reported thus far [70,83,84,98]. However, a reasonable reaction scheme can be formulated if all the known evidence is viewed collectively. This evidence will be briefly summarized at the outset. First, control experiments indicate no diamond growth either in the absence of laser excitation or when CH4 is absent from the gas mixture. Further, no diamond growth was obtained from optically pumped CO/C2H2 mixtures. Second, elemental oxygen is detected in the EDS spectra of diamond particles grown using LECVD as well as laser photolysis. This suggests that CO somehow acts in concert with CH4 during growth. Third, there is clear evidence of energy transfer from CO to CH4 as can be seen from the IR emission from excited states of CH4. Fourth, the LECVD experiments conducted at E´cole Centrale Paris detected only CO and CH4 and no other species (within detection limits) in their emission and absorption studies [98]. Fifth, the morphology of the irregularly faceted particles grown by LECVD is identical to those grown by laser photolysis. Finally, the particle growth rates appear to be insensitive to whether or not the substrate is heated. Based upon these six essential points, we attempt to formulate a possible mechanism. There have been several important attempts to elucidate the surface mechanisms responsible for diamond growth [64,65,99–102]. These mechanisms are relevant for conventional CVD techniques for diamond growth in a hydrogen-rich environment and have shown promise in providing a mechanistic understanding of growth on (100) surfaces. Many of the proposed mechanisms focus on the incorporation of CHx species on the surface, specifically the CH3 molecule (i.e., methyl radical) [64,65,99–101]. We therefore concentrate here on the possibility of extending these ideas to understand the LECVD process. However, it must be borne in mind that
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these mechanisms have been postulated for a hydrogen-rich ambient and not for the conditions typical of LECVD. The first point that must be addressed is the production of methyl radicals either in the gas phase or on the surface of seed diamond particles. Given the relatively high vibrational temperature of CO (⬃2200 K) produced by laser excitation and subsequent VV pumping, it is reasonable to expect the translational temperature of CH4 to be between 300 and 2200 K. Under these conditions, production of CH3 can occur via the following channels: Vibrationally stimulated thermal dissociation: CH4 ⫹ M → CH3 ⫹ H ⫹ M Vibrational pooling:
CH* 4 ⫹ CH* 4 → CH3 ⫹ H ⫹ CH4
(32) (33)
CO(v) ⫹ CH* 4 → [CH3COH]* → [C2H4O]* → [CH3CHO]*
CH3 ⫹ HCO CH3CHO CH4 ⫹ CO
(34)
where the * denotes an excited species. The formation of [CH3COH]* can proceed via an insertion reaction whereby a CO molecule inserts itself into a vibrationally excited CH4. Such insertion reactions involving AlO inserted in CH4, have been previously reported [103]. The last step of process (34) has been observed in studies of decomposition of oxirane [104]. It is also interesting to note that CH3CHO is more stable than C2H4O by about 27 kcal/mol or ⬃1.17 eV/molecule. But, there is an activation energy of ⬃57 kcal/mol or ⬃2.47 eV/molecule for the isomerization C2H4O → CH3CHO. This activation energy can be provided via the vibrational excitation of the reactant CO and CH4 molecules by pooling because energies up to ⬃2 eV are available from CO molecules near v = 10 (see Fig. 45). A compelling argument can be made for processes (33) and (34) based upon energy transfer considerations. It can be seen from Fig. 44 that the 3 mode of CH4 cannot be populated by intermolecular energy transfer from CO(v = 1) because its energy is higher than the energy corresponding to the v = 1 → v = 0 transition.24 However, intermolecular energy transfer can occur to the 2 and 4 states of CH4 via CO(X1⌺⫹, v = 1) ⫹ CH4 → CO(X1⌺⫹, v = 0) ⫹ CH4(2 or 4)
(35)
It can thus be seen that the 3 mode can be populated by VV up-pumping 24
Two-quantum transfer via processes such as CO(v = 2) ⫹ CH4 → CO(v = 0) ⫹ CH4(3) are possible but much less probable than transfer of single vibrational quanta.
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Table 6 Typical Characteristics of Low-Lying Energy Levels and VV Transfer in Methane at 2.5 Torr State
1 (symmetric stretch) 2 (symmetric bend) 4 (asymmetric bend) 3 (asymmetric stretch) 22 24 23 23 → 3 2 ⫹ 4 3 → 4 3 → 2 1 → 4 24 → 2 ⫹ 4 1 ⫹ 4 → 3
Characteristic time for VV transfer
Radiative lifetime
O(0.1–0.5 sec) O(0.1–0.5 sec) 1.1 sec 1.6 sec O(0.1–0.5 sec) O(0.1–0.5 sec) O(0.1–0.5 sec) 0.92 sec O(0.1–0.5 sec) 0.72 sec O(0.1–0.5 sec) O(0.1–0.5 sec) 1.6 sec 0.22 sec
>20 sec 13 sec 0.39 sec 37 msec — 0.20 sec — — 0.39 sec 25 sec 120 sec 8 sec — —
Source: Refs. 105–107.
within the methane initiated from either the 2 or 4 modes25 or by radiative decay from an excited state. In the latter case, an upper state could have been populated only by VV up-pumping within the methane. Once a specific excited state is populated, it can undergo one of several processes. It can radiate a photon (provided selection rules allow such a transition) and occupy a lower energy state or ground state. It can be collisionally deexcited to any lower state by a VT-type collision or collisionally excited to an upper state by intermolecular or intramolecular vibrational energy transfer. In CH4, the radiative lifetimes of the low-lying levels are longer than tens of milliseconds, as can be seen from Table 6 [105–107]. In contrast, VV transfer occurs within microseconds at pressures of interest to LECVD (i.e., PCH4 = 2.5 torr) [106]. Also, VT and VR collisional deactivation of these levels occurs on time scales of hundreds of microseconds or milliseconds at 2.5
25
The 2 and 4 levels of methane are isoenergetic with the v = 24 → v = 23 and v = 33 → v = 32 transitions in CO, respectively. Because such high vibrational levels are not significantly populated in CO (as can be seen from the overtone emission spectrum of CO), these methane energy levels must have been populated by off-resonance intermolecular energy transfer from v → v ⫺ 1 transitions in CO (with the v = 1 → v = 0 transition being the most likely one).
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torr [106]. Hence, if intermolecular energy transfer from CH4 back to CO does not depopulate the upper levels of CH4 appreciably, a compelling argument can be made for rapid VV exchange among CH4 molecules. We therefore expect CH4 molecules to be vibrationally hot. This is evident from the strong radiative emission from the 3 mode of CH4 despite the presence of rapid intramolecular VV transfer and self-absorption of radiation from the 3 mode by surrounding, cold CH4. It is therefore clear that vibrational energy from CO is rapidly transferred to CH4, thereby populating the 2 or 4 (or both) states. The presence of vibrationally excited CO and excited CH4 in the LECVD experiments has been inferred. These can lead to formation of CH3 radicals via processes (32) through (34). Similar reactions could occur between CO and CH4 molecules adsorbed on diamond seed surfaces as well, and there is not enough experimental evidence to argue for production of methyl radicals from the gas phase alone. Nevertheless, given such an initial source of production of CH3 radicals, the evidence gathered from the LECVD experiments collectively leads to two possible mechanisms. They are (1) growth from CH3 radicals with CO playing no part other than that of a spoiler and terminating growth and (2) growth by concerted action of CO and CH3 at the surface of the diamond seed particles. Let us consider these two possible mechanisms in some detail, especially from the perspective of experimental observations. For ease of discussion, we consider growth on the (100)-oriented surface. Once CH3 radicals are produced (via either gas-phase or surface processes), they can attach themselves to dangling bonds on the (100) surface as shown in Fig. 49. This can be written symbolically as: C•d ⫹ CH3 → Cd — CH3
(36)
• d
where C represents an activated site (i.e., dangling bond) on a seed diamond surface. Atomic hydrogen from the CH3 group can then be removed by successive attacks by other CH3 radicals: Cd — CH3 ⫹ CH3 → Cd — CH2• ⫹ CH4(g)
(37)
In this manner, carbon can be successively added to and atomic hydrogen
Figure 49 Schematic showing attachment of CH3 radical on a dangling bond on the (100) surface of a diamond speed particle.
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removed from the surface by CH3 radicals alone. Molecular dynamics calculations have shown that a high fraction of CH3 radicals (more than 0.5 at normal incidence) can be adsorbed on (100) diamond surfaces over kinetic energy ranges from 0.5 to 2 eV [108]. Other processes for abstraction of H atoms can also take place. One such process occurs by attack from vibrationally excited CO as suggested previously [109]: Cd — CH3 ⫹ CO() → Cd — CH2 ⫹ HCO
(38)
The other can occur by the following process (after abstraction of an H atom from each of the CH3 groups): Cd — CH2• ⫹ Cd — CH2• → Cd — C2H4
(39)
In this scenario, the CO serves the sole purpose of exciting CH4 and producing CH3. It can also play the role of a spoiler by occupying a dangling bond and establishing a ketone-type structure. This structure is stable (unless attacked by a radical such as CH3) and can in principle inhibit growth by the CH3 mechanism just discussed. The ketone and other possible surface structures are shown in Fig. 50. The second possible mechanism requires the presence of both CO and CH4. Here, CO is added to the surface (the activation energy for adsorption supplied by vibrational excitation) so as to form a ketone-type structure (see Fig. 50a): —O C•d ⫹ CO → Cd — C —
(40)
The adsorbed CO is then attacked by a methyl radical, resulting in the conversion of the double bond between C and O to a single bond by scission:
Figure 50 Schematic showing attachment of CO on the (100) surface of a diamond seed particle, forming a stable ketone-type structure (a) and an ether-type structure (b).
Laser-Assisted and Optical Pumping Techniques
— O ⫹ CH3 → Cd — C — O — CH3 Cd — C —
407
(41)
The resulting ether-like structure is shown in Fig. 51. After subsequent Hatom abstraction via (37), (38), or (39), an aldehyde-type structure can form momentarily, leading to -scission of the underlying CO bond. This should result in evolution of formaldehyde (H2CO), leaving behind two dangling bonds (i.e., two active sites): Cd — C• — O — CH2• → Cd — C• ⫹ H2CO(g) The two dangling bonds on the carbon now allow another layer to be grown by repetition of the preceding processes. Both mechanisms serve to explain some of the key observations in the LECVD experiments. Both can account for the presence of small amounts of elemental oxygen (see the EDS spectrum shown in Fig. 42), and both give plausible descriptions of growth in an environment depleted of atomic hydrogen. Of the two, however, the first appears to explain additional observations reported in Ref. 98. The result of hydrogen abstraction by methyl radicals in the first mechanism would result in the evolution of CH4 from the surface, which would be indistinguishable from the methane already present in the gas phase. This would explain why only CO and CH4 were detected in the absorption measurements reported in Ref. 98. The question remains why CH3 radicals or perhaps HCO went undetected. A possible answer to this question is that CH3 concentration levels (according to our estimates, methyl radical concentrations on the order of nCH3 ⬇ 1012 cm⫺3 or 1018 m⫺3 are expected, or about 1 ppm) were well below detection limits. By this latter argument, the second mechanism cannot be ruled out because
Figure 51 Schematic showing the result of attack by the CH3 radical on a ketonetype CO structure occupying a dangling bond on the (100) surface of a diamond seed particle.
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evolution of H2CO and HCO levels could have been below detection limits as well. Unfortunately, ultimate quantitative detection limits were not reported in Ref. 98 so that it is not possible to carry out further analysis of possible mechanisms. Given the current level of interest in synthesizing diamond at low substrate temperatures [76], both LCVD and LECVD appear to be significant advancements toward attaining this goal. Both laser-based techniques challenge the conventional myth that high substrate temperatures are necessary for diamond growth. Clearly, the environment of conventional diamond growth techniques is limited only by how energy is loaded into the gas mixture. Both techniques also offer some selectivity as can be seen by the conspicuous lack of codeposition of nondiamond carbon, prevalent in other CVD methods. LECVD also challenges our understanding of diamond growth mechanisms by displaying growth in an environment that is not rich in atomic hydrogen. Ultimately, that may be the reason why growth rates in the two methods are small—on the order of 1 m/hr. For LECVD, as the absorbed power is on the order of 1 W, we may estimate the energy cost per carat as 1 J/sec ⫻ (1/10⫺12 cm3) ⫻ (1 hr ⫻ 3600 sec/hr) ⫻ (1/3.515) (cm3/g) ⫻ (0.2 g/carat) = 2 ⫻ 1014 J/carat. This value is many orders of magnitude higher than the corresponding value for diamond synthesis by an oxyacetylene flame. However, unlike LCVD and other methods of CVD diamond synthesis, the LECVD process scales with pressure and can be operated potentially at pressures on the order of 20 atm [77,83], so that the J/carat value can be reduced drastically. However, continuous films have yet to be produced by either LCVD or LECVD. The optical pumping technique shows tremendous promise for synthesis of other materials. For example, it has been shown that vibrationally excited nitrogen is important in the production of silicon nitride films [110]. Reference 110 reports that when silane was added to a flowing nitrogen discharge afterglow (where there is known to be copious amounts of vibrationally excited N2 molecules) and the reaction products made to impinge on a heated monocrystalline silicon substrate placed downstream, silicon nitride films were produced. In addition, reactions between the vibrationally excited N2 and silane in the afterglow led to formation of fine (micron-sized) particles of white silicon nitride powder [110]. Given the similarity in structure between CH4 and SiH4, it is likely that optically pumped CO/SiH4 mixtures could be used for deposition of silicon at low temperatures or for synthesis of silicon carbide. Indeed, Ref. 110 determined experimentally that about 15% of the vibrationally excited N2 quenching events resulted in excitation of the 3 mode of SiH4. Recently, a combination of optical pumping and application of a subbreakdown radio-frequency electric field has been shown to produce plasmas in CO/N2 gas mixtures seeded either with NO or
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O2 at total pressures ranging from 0.4 to 1.2 atm [111]. Such plasmas, although weakly ionized, contain substantial amounts of vibrationally excited CO and N2. The use of such plasmas for synthesis of novel materials remains to be explored. In this section, we have examined a laser excitation technique known as LECVD for diamond synthesis. Of significance is the fact that growth of diamond particles has been realized under conditions where the substrate is not heated at all [70,84] and without the need to hyperdilute the gas phase with hydrogen. In fact, diluents such as hydrogen poison the vibrational uppumping of CO and CH4, an effect that is central to the LECVD process. In the following section, we examine a potential means of extending the LECVD process to the condensed phase in order to increase growth rates. C.
Optical Pumping of Liquids
Energy transfer as well as the optical pumping of gases was discussed in Sec. V.B. Specifically, it was shown that small polyatomic molecules are very effective in retaining energy in their vibrational modes. For diatomic molecules, in the limit as VV exchange rates dominate VT and VR energy transfer rates, the population of various vibrational levels is described by the TRR distribution for anharmonic oscillators [see Eq. (17)]. As long as VV rates exceed VT, VR, and spontaneous emission rates, energy is effectively retained in the vibrational modes of molecules. It is important to point out that vibrational energy exchange among anharmonic oscillators is a collisional process. As long as the energy defect26 is carried away by a suitable cold bath, energy can be maintained for long times within the vibrational modes. This is irrespective of the pressure or density of the medium. In the gas phase, CO/Ar mixtures have been optically pumped in a manner so as to excite CO to very high vibrational states (near v = 40) at pressures of 20 atm [77]. Such strong vibrational excitation has also been observed in the liquid and solid phases. This is of significance if we are interested in driving chemical reactions, while maintaining some degree of control and selectivity over which reactions occur and to what extent. In this section, we describe extension of the LECVD technique described in Sec. V.A and V.B to the condensed phase. 26
Recall that the energy defect is defined as the energy difference between the total energy of the colliding molecules before the collision and that value after the collision. During anharmonic VV pumping, energy of magnitude equal to the energy defect is transferred from vibrational excitation to the external modes of translation and rotation. This results in eventual heating of the gas unless the gas is flowed or a suitable diluent is used to carry off this excess heat.
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There exists a considerable literature on optical pumping of CO at elevated gas pressures as well as in the liquid and solid phases [77,112– 118]. These studies have been largely motivated by interest in fundamental energy transfer processes, isotope separation processes, and development of energy transfer lasers. There exists a relatively good coincidence between the laser lines emanating from a CO laser operating on the abundant isotope 12 16 C O and the isotopic liquid phase mixture of 88% 13C16O and 12% 13C18O. Several researchers have used this coincidence to excite CO vibrationally in the liquid phase and in the solid phase [112–118]. Remarkably, the condensed phase CO behaves as though it were a dense gas within the liquid argon matrix or on the solid single-crystal NaCl (100) surface [112–117]. Processes encountered in the gas phase (as discussed in Sec. V.A. and V.B) also occur in the condensed phases, and very high vibrational levels near v = 40 in CO have been observed [112]. Further, energy transfer from CO to small polyatomic molecules has also been reported [112,117,119,120], and not surprisingly, vibrationally excited CO and O2 have been reacted to form CO2 in the liquid phase at temperatures near 88 K. This is indeed a dramatic demonstration of the extreme disequilibrium between vibrational modes and external modes created by optical pumping. Although diamond growth from such optically pumped CO/CH4 liquid mixtures has never been demonstrated, the successful LECVD experiments described in Sec. V.B are strong indicators that diamond growth in optically pumped liquids could meet with success. Further, measured rate coefficients for energy transfer in the liquid and gas phases for a variety of mixtures of interest to diamond synthesis are quite encouraging. Table 7 gives measured values of rate coefficients for gas and liquid phases at a temperature of 80 K, as reported [119,120]. Remarkably, the variation in rate coefficients for vibrational energy transfer between gas and liquid states is less than 10% for the molecules listed in Table 7. Now, the rate of energy transfer depends on the product of the rate coefficient and number densities of the colliding partners. Therefore, the rate at which intermolecular vibrational energy transfer occurs increases by many orders of magnitude because the liquid phase is much denser than the gas phase. This has the potential to increase yield and growth rates. As many single-crystal materials are grown from melts in the liquid phase, optical pumping in the liquid phase may finally lead to epitaxial, single-crystal diamond films. The latter possibility is further bolstered by the fact that multilayers of ␣-form solid, single-crystal CO can be deposited on NaCl (100) substrates epitaxially because the lattice mismatch is less than 0.5% [116]. Further, the solid CO has been vibrationally excited up to levels as high as v = 23 [114], and as high as v = 30 at 22 K, by optical pumping using a CO laser [116]. In fact, this up-pumping is so rapid in the solid phase as to yield a population inversion (i.e., there are greater
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Table 7 Measured Rate Constants for Vibrational Energy Transfer in the Gas and Liquid Phases at 85 K System N2-12C16O N2-13C16O N2-13C18O 13 18 C O-O2 13 16 C O-O2 12 16 C O-O2 13 18 C O-CH4 13 16 C O-CH4 12 16 C O-CH4 N2-CH4 12 16 C O-CH4 13 16 C O-CH4 13 18 C O-CH4 12 16 C O-CF4 13 18 C O-CF4 12 16 C O-CD4 13 16 C O-CD4 13 18 C O-CD4 12 16 C O-CD3H 13 16 C O-CD3H 13 18 C O-CD3H
⌬E (cm⫺1) 187 234 286 487 540 587 509 562 609 796
kLiquid (cm3/sec) (3.2 ⫾ 0.2) (1.2 ⫾ 0.1) (3.9 ⫾ 0.1) (1.54 ⫾ 0.02) (7.0 ⫾ 0.5) (4.2 ⫾ 0.1) (4.03 ⫾ 0.09) (2.08 ⫾ 0.05) (1.33 ⫾ 0.05) (1.3 ⫾ 0.1) (1.33 ⫾ 0.05) (2.08 ⫾ 0.05) (4.03 ⫾ 0.09) (8.6 ⫾ 0.1) (7.50 ⫾ 0.09) (2.48 ⫾ 0.04) (2.28 ⫾ 0.08) (1.47 ⫾ 0.03) (1.15 ⫾ 0.02) (1.58 ⫾ 0.04) (1.26 ⫾ 0.02)
⫻ ⫻ ⫻ ⫻ ⫻ ⫻ ⫻ ⫻ ⫻ ⫻ ⫻ ⫻ ⫻ ⫻ ⫻ ⫻ ⫻ ⫻ ⫻ ⫻ ⫻
kGas (cm3/sec) 10⫺15 10⫺15 10⫺16 10⫺17 10⫺18 10⫺18 10⫺15 10⫺15 10⫺15 10⫺16 10⫺15 10⫺15 10⫺15 10⫺14 10⫺16 10⫺13 10⫺13 10⫺13 10⫺12 10⫺13 10⫺13
(3.6 ⫾ 0.2) (1.1 ⫾ 0.1) (3.79 ⫾ 0.08) (1.50 ⫾ 0.03) (7.0 ⫾ 0.9) (4.1 ⫾ 0.2) (4.17 ⫾ 0.08) (2.07 ⫾ 0.06) (1.27 ⫾ 0.04) (1.5 ⫾ 0.2) (1.27 ⫾ 0.04) (2.07 ⫾ 0.06) (4.17 ⫾ 0.08) (1.98 ⫾ 0.04) (1.0 ⫾ 0.1) (4.60 ⫾ 0.08) (4.3 ⫾ 0.1) (3.5 ⫾ 0.1) (1.32 ⫾ 0.03) (3.3 ⫾ 0.1) (2.38 ⫾ 0.05)
⫻ ⫻ ⫻ ⫻ ⫻ ⫻ ⫻ ⫻ ⫻ ⫻ ⫻ ⫻ ⫻ ⫻ ⫻ ⫻ ⫻ ⫻ ⫻ ⫻ ⫻
kLiquid/kGas 10⫺15 10⫺15 10⫺16 10⫺17 10⫺18 10⫺18 10⫺15 10⫺15 10⫺15 10⫺16 10⫺15 10⫺15 10⫺15 10⫺13 10⫺15 10⫺13 10⫺13 10⫺13 10⫺12 10⫺13 10⫺13
0.97 1.09 1.03 1.03 1.00 1.02 0.97 1.00 1.05 0.87 1.05 1.00 0.97 0.43 0.75 0.54 0.53 0.42 0.87 0.48 0.53
⫾ ⫾ ⫾ ⫾ ⫾ ⫾ ⫾ ⫾ ⫾ ⫾ ⫾ ⫾ ⫾ ⫾ ⫾ ⫾ ⫾ ⫾ ⫾ ⫾ ⫾
0.08 0.13 0.03 0.03 0.15 0.06 0.03 0.04 0.06 0.13 0.06 0.04 0.03 0.01 0.08 0.01 0.02 0.02 0.02 0.02 0.01
Source: Refs. 119, 120.
numbers of molecules in the higher vibrational states than in the v = 0 or ground vibrational state). We now describe an experiment that demonstrates some of the salient features of optical pumping of CO in the liquid phase [70]. A CO laser designed to operate cw (5 to 8 W on all lines) with some laser output on the v = 1 → v = 0 transition was used to pump CO in a liquid argon matrix. A schematic of the cell containing the cryogenic mixture is shown in Fig. 52. This cell is identical to that used in the experiments reported in Ref. 112, and details may be found therein. The output from the CO laser was focused into the center of the cryogenic cell containing a solution of CO in liquid argon, using a ZnSe lens (1 in. diameter, focal length of 9 in.). The liquid mixture was produced from an initially gaseous mixture of CO and argon, which was condensed upon being cooled by the liquid argon placed in a dewar immediately above the cell. A schematic of the mixing chamber
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Figure 52 Schematic of the cell containing the cryogenic mixture of CO and argon. (From Ref. 70.)
is shown in Fig. 53. The CO/argon mixture is prepurified before allowing it to condense into the cell. The purification procedure consists of using the cold finger connected to the mixing chamber. The gases are condensed into the cold finger and, after some time, allowed to evaporate back into the mixing chamber while a small part of the cold finger is still immersed in liquid nitrogen. This process is repeated 15 times. Reference 70 reported pumping of CO/argon gas mixtures where the initial partial pressure of CO in the gas phase was varied between 70 mtorr and 30 torr, while maintaining a total pressure of 2 atm. The CO was an isotopic mixture comprising 88% 13C16O and 12% 13C18O. This isotopic mixture was selected because of its excellent coincidence in the liquid phase with CO laser lines when the CO laser was operated on the normally abundant 12C16O isotope. Following the aforementioned purification procedure, the CO/argon mixture was condensed into the cell and irradiated by the cw CO laser beam. Power measurements before the cell were in the range of 3.5 to 5 W. Using a side port, IR emission was monitored from the optically pumped liquid CO. A representative spectrum of the first overtone (⌬v = 2)
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Figure 53 Schematic showing the mixing chamber used in optical pumping of liquid CO/argon mixture. (From Ref. 70.)
transitions from the liquid CO for an initial gas-phase CO partial pressure of 9 torr, after background blackbody subtraction, is shown in Fig. 54. It is important to note that this is not a low-resolution spectrum. Rotational lines within each vibrational band can be either totally absent (as in the case of CH4) or substantially narrowed in the liquid state as compared with similar gas-phase spectra. It is immediately apparent from the emission spectrum that states at least as high as v = 38 (see the energy level diagram for CO
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Figure 54 A representative spectrum of the first overtone (⌬v = 2) transitions from liquid CO in liquid argon, optically pumped using a CO laser. This spectrum was obtained for an initial gas-phase CO partial pressure of 9 torr and after background blackbody subtraction. (From Ref. 70.)
shown in Fig. 34) of the heavier isotope 13C18O are substantially populated. In addition, peaks corresponding to isotopes of carbon dioxide 13C16O16O, 13 16 18 C O O, and 13C18O18O are also observed. This is indicative of chemical reaction occurring among vibrationally excited CO molecules, possibly via channels (29). Production of CO2 has been observed previously but from vibrationally excited CO doped with O2 [112]. Assuming the Einstein A coefficients for spontaneous emission from these vibrational states are comparable between gas and liquid phases, relative populations can be inferred from the emission spectrum. This is shown in Fig. 55, displaying a fully VV pumped distribution in the heavier isotope. It is clear that many (if not all) of the processes observed in gas-phase optically pumped CO occur in the liquid phase as well. Anharmonic VV pumping resulting in highly vibrationally excited CO (levels as high as v = 38 or as energetic as 8 or 9 eV) is observed in both gas and liquid phases. Intermolecular energy transfer between CO and CH4 has been reported, although IR emission from highly excited methane has not yet been observed. Chemical reaction from vibrationally excited CO, and between CO and O2, yielding CO2 is observed in both the gas and liquid phases as well. Given this close similarity in the occurrence of key processes between gas and liquid phases and the success of the LECVD process for diamond synthesis, the likelihood that diamond can be synthesized from optically pumped liquid CO/CH4 /Ar mixture appears to be high. Should diamond synthesis occur in
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Figure 55 Relative populations of vibrationally excited CO in liquid argon inferred from the emission spectrum. This was obtained by assuming that the Einstein A coefficients for spontaneous emission from these vibrational states are comparable between gas and liquid phases. Note the fully VV pumped distribution in the heavier isotope. (From Ref. 70.)
such an optically pumped liquid mixture, it should be more rapid than in the gas phase because the rates of energy transfer scale as the square of the density. Finally, the production of flaw-free, cubic, solid CO as well as its extreme vibrational excitation by optical pumping resulting in population inversion bodes well for adsorbed CH4 species. Although growth of diamond films using such a technique has not been demonstrated, it is nevertheless worth exploring. Given the fact that many growth techniques for single crystals of other materials involve the liquid phase, this approach of optically pumping liquids appears promising for synthesis of single-crystal diamond. This is important before diamond can be considered as a candidate material for electronic applications.
VI.
SUMMARY
This chapter has discussed synthesis of diamond from gas and liquid phases using laser sources. The laser-based techniques include surface heating, py-
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rolysis, ablation (LPVD), photolysis (LCVD), and optical pumping (LECVD). The laser-based techniques do not necessarily supplant existing CVD methods but offer the prospect of lowering growth temperatures, selectively depositing diamond while inhibiting the deposition of nondiamond carbon, and perhaps enhancing growth rates. In the case of rapid gas-phase heating such as pyrolysis, the laser simply serves as a source of thermal heating and is therefore an inefficient means of activating the gas. In this instance, it does not provide any advantages such as selectivity or enhanced quality of the deposit. Ablation and heating of surfaces can benefit from laser sources because reactions may need to be driven locally, without bulk heating of the growth medium, target, or substrate. In contrast, photolysis and optical pumping techniques use attributes of molecular structure to drive specific reaction pathways to form products. This is difficult or not even possible using conventional growth methods. Many of these laser-based processes have already demonstrated diamond growth, albeit at different rates. Although capital costs vary widely, a useful measure of the relative efficacy of each process is the joule-per-carat value. This measure considers the power required for the process as well as growth rate simultaneously. The pyrolysis and photolysis techniques exhibit growth rates on the order of microns per hour, as does the optical pumping technique at low gas pressures. However these exhibit drastically different energy costs in J/carat: O(108), O(107)–O(1014), and O(1014), respectively. These should be contrasted with typical values of O(108) J/carat for HFCVD27 and O(106) J/carat for synthesis using an oxyacetylene flame.28 The laser-plasma technique is equivalent to but less efficient energetically than the conventional oxyacetylene flame, exhibiting growth rates on the order of 10 to 100 m per hour at a power level of 2.5 kW. Of all these laser processes, however, the well-publicized but as yet unverified QQC process reports growth rates on the order of millimeters per hour (i.e., microns per second) at a cost of about 71 J/carat. Clearly, if cost is the prevailing factor, the QQC process appears to be the best (provided its claims are substantiated), while some of the laser pyrolysis and laser photolysis methods are competitive with HFCVD and oxyacetylene flame synthesis. It would be premature to eliminate optical pumping based on a comparative measure such as joules per carat. Optical pumping, unlike the other methods, inherently scales directly with pressure. Therefore, for approximately the same power expenditure, the growth rate can be significantly increased sim-
27
This value is based on a power expenditure of 100 W and deposition over a 1 cm2 area at 1 m/hr. 28 See footnote 9.
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ply by raising the pressure. This would reduce the J/carat value for this process from the unreasonable O(1014) and make it competitive with the others. Further enhancement in growth rate is possible in the liquid phase, in principle, although diamond growth in an optically pumped liquid is yet to be demonstrated. Moreover, this method has the added virtue of lower growth temperatures as well as process selectivity. Some of the more nascent but novel approaches such as LECVD, laser photolysis, and the QQC process have been critically examined. Some appear promising in terms of scaling to higher growth rates, lower growth temperatures, and providing selectivity. Conventional CVD and some of the laser-based (such as pyrolysis and laser-plasma) methods are useful for substrates that can withstand high temperatures and the presence of large amounts of atomic hydrogen or oxygen. The more novel of the laser-based methods show some promise of being able to circumvent these difficulties. It is particularly noteworthy that the optical pumping technique is capable of activating gases over a wide a range of pressures, as well as liquids. There are precedents for growth of single crystals of various materials from the liquid phase. Optical pumping lays the groundwork for development of a new class of processes that employ laser excitation to drive energy transfer and chemistry in cryogenic liquids and may someday play a role in the production of single-crystal diamond films.
ACKNOWLEDGMENTS Partial support from grants MSS-9157303 and MSS-9102076 from the National Science Foundation, The Ohio State University College of Engineering, The Center for Materials Research, The Ohio Board of Regents Investment Fund Program, and the Directorate of Defense Research and Engineering (DDR&E) under the Air Plasma Ramparts MURI program managed by the Air Force Office of Scientific Research is gratefully acknowledged. The authors thank Dr. Gregory Hall, Chemistry Department, Brookhaven National Laboratory for many helpful comments and suggestions.
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10 CVD Diamond Solutions for Machining and Other Mechanical Applications Brian L. Cline Cline Innovations, LLC, Sterling, Massachusetts
James M. Olson Fairchild Semiconductor, South Portland, Maine
I. A.
INTRODUCTION Growth Engine for Advanced Cutting Tool Materials
The machining industry is always seeking a means of improving productivity and machined part quality. Over the years, this relentless demand has resulted in countless iterations of productivity improvements. Today, the machine tool industry produces automated high-speed machining centers that have the ability to machine complex shapes quickly with high accuracy and minimal operator assistance. These dramatic machine tool improvements create an interrelated demand for improved cutting tool materials that enhance productivity by extending tool life under increasingly aggressive machining conditions. Improvements to cutting tool materials are also needed for enhanced cutting tool performance in machining many advanced workpiece materials. For example, the aerospace and automotive industries are utilizing increasing quantities of lightweight materials such as high-silicon aluminum alloys, fiber-reinforced composites, and metal matrix composites (MMCs), which offer unique options for component design. Unfortunately, many of these enabling materials are very abrasive and difficult—if not impossible—to machine using conventional cutting tools [1–3]. 425
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This chapter contains an overview of how and why chemical vapor deposition (CVD) diamond technology is beginning to assume a key role in the materials evolution for advanced cutting tool and specialty wear part applications. Section I continues with a history of cutting tool materials development including high-speed steel, cemented carbides, hard coatings, and synthetic diamond materials. Section II overviews the issues and challenges associated with CVD diamond tool design and manufacture. This section also includes discussion of CVD diamond properties and application specifics that can affect the performance of CVD diamond cutting tools as well as common approaches to CVD diamond tool design. Section III is focused on microstructural features of CVD diamond that affect the performance of diamond coatings as well as freestanding diamond in mechanical applications. Section IV contains cutting performance information for various CVD diamond tool types used in a wide range of materials. The chapter closes with an overview of other types of mechanical component applications in Sec. V followed by a summary in Sec. VI. B.
Evolution of Conventional Cutting Tool Materials
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Carbon Tool Steels and Nonferrous Cast Alloys
Early industrial tooling materials included carbon tool steels—iron-based materials with a wear-resistant ‘‘case’’ surrounding a tougher FeC alloy. Although these materials display high fracture resistance, their low hardness and low wear resistance limit their wear life in industrial machining of metals and highly abrasive materials. The introduction of high-speed steels (containing varying amounts of W, Cr, V, Mo, and Co) led to the rapid replacement of carbon tool steels in metalcutting applications due to marked improvements in wear and fracture resistance. The added wear resistance and high-temperature hardness properties (‘‘hot hardness’’) of nonferrous cast alloys (containing varying amounts of Cr, W, C, and Co) allowed further improvements in cutting capabilities for selected metalcutting operations. 2.
Cemented Carbides Including Tungsten Carbide
The subsequent development of cemented carbides in the late 1920s provided an immediate advantage over traditional high-speed steels and cast alloys [4]. The combination of properties offered by cemented carbide cutting tools is typically attained by sintering large volumes of ceramic grains with 3 to 15% (by volume) of metallic binder and various sintering aids. Cobalt-sintered tungsten carbide combines wear-resistant, hard tungsten carbide (WC) ceramic particles with the fracture resistance of a ductile, metallic, Co-based binder phase. Many ‘‘tungsten carbide’’ grade formulations are
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readily available for a wide range of applications in the cutting tool and wear part industry. The metal sintering approach to engineering cutting tool materials bypasses the challenges of introducing the high hardness and wear resistance of otherwise brittle ceramic materials into cutting tools. Although WC-Co cutting tools have lower fracture resistance than high-speed steel, their mechanical properties are suitable for a very wide range of heavy-duty applications. The higher hardness and improved hot hardness characteristics of cemented carbides allow operation at cutting speeds three to four times higher than with high-speed steel. C.
The Historic Use of Diamond and Ceramic Cutting Tool Materials
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Single-Crystal Diamond
For decades, natural and high-pressure, high-temperature (HPHT) singlecrystal diamonds have been used for specialty cutting and wear applications. Single-crystal diamond tools are known to have unmatched edge quality and create excellent surface finish due to their precision, continuous diamond edges. Applications include woodworking, contact lens machining, jewelry manufacture, and nonferrous metal finishing. The cost and difficulty of fabricating tools of complex geometries limit the usage of single-crystal tips to simple shapes relative to other tooling materials. However, the hardness, thermal stability, and corrosion resistance of single-crystal diamond tooling continue to provide added value in many demanding industrial applications. Single-crystal tools are also limited by their low fracture resistance, anisotropic wear behavior, and morphological shape constraints relative to other forms of diamond. 2.
Polycrystalline Diamond Compacts
In the early 1970s, technology was developed that allowed the production of metal-sintered polycrystalline diamond (PCD) compacts in large, twodimensional flat plates sintered directly to a cemented carbide backing. Similar in construction to cemented carbide, PCD diamond tooling has significant advantages related to the high wear resistance of diamond grains combined with the fracture resistance afforded by the metallic binders used in sintering. Figure 1 contains scanning electron microscope (SEM) images that compare the two-phase microstructures of a cemented carbide cutting tool and a PCD cutting tool. The added hardness and wear resistance of the diamond grains in PCD relative to the tungsten carbide grains in WC-Co offer distinct advantages for PCD-tipped tooling. PCD tools are normally fabricated by radio-frequency or torch brazing flat, tungsten carbide-backed,
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polycrystalline tool tips directly onto a cemented carbide or high-speed steel tool body. After brazing in air, PCD tools are typically finish ground or cut to final shape by electric discharge machining (EDM). The performance advantages of PCD over tungsten carbide tools clearly outweigh the added cost of brazed PCD tooling in many nonferrous metal and abrasive nonmetallic machining applications. PCD is commonly used for machining aluminum, copper, brass, bronze, and cemented carbide in applications including turning, boring, profiling, grooving, milling, and holemaking [5]. 3.
Advanced Ceramics
Ceramic cutting tools are produced from a wide range of polycrystalline ceramic materials including aluminum oxide, cubic boron nitride (‘‘PCBN’’), and silicon nitride [6–9]. The hardness, chemical stability, and high wear resistance of these advanced ceramic cutting tools have demonstrated costperformance advantages over conventional cutting tool materials in applications ranging from abrasive and corrosive polymers to superalloy machining. PCBN-tipped tools are rapidly becoming the favored cutting tool choice in many high-throughput ferrous applications. Although the wear resistance, high-temperature corrosion resistance, and hot hardness of ceramic cutting tools are impressive, the application range of ceramic tooling is often limited by the lower fracture resistance of ceramic materials relative to other cutting tool materials. Special cutting edge preparation, whisker reinforcement, or other approaches are commonly used to improve the performance and broaden the application range of ceramic cutting tools. D.
Development of Coating Technologies for Cutting Tools
In the 1970s, thin-film deposition technology began to play a key role in furthering performance of commercial cutting tools [10]. Coated carbide and
< Figure 1 Structural similarities between cemented tungsten carbide and PCD diamond: SEM photomicrographs showing (a) the fractured cross section of a 6% CoWC cemented carbide at 5000⫻ magnification, (b) a polished cross section of a 6% Co-WC cemented carbide at 2000⫻ magnification, and (c) a 50⫻ image showing the construction of a PCD cobalt-sintered diamond cutting tool. Because both cemented carbides and PCD are two-phase materials composed of hard ceramic grains (WC or diamond, respectively) sintered in metal, they are commercially available in various compositions and grain sizes tailored for a wide range of applications. (From Norton Diamond Film/Saint Gobain Industrial Ceramics, Inc.)
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coated high-speed steel tools utilize hard coatings such as titanium nitride, titanium carbide, and aluminum oxide to enhance cutting performance. Most coatings are produced using CVD or physical vapor deposition (PVD) techniques. This ‘‘layer-composite’’ approach to cutting tool design couples the high-temperature wear resistance, corrosion resistance, and hardness advantages of advanced ceramic materials with the superior fracture resistance of the underlying tool body. Today, hard coating technology has evolved to the development of multilayered coatings in which individual coating layers offer unique performance-enhancing properties beyond single-layer structures [10,11]. Coating technologies for tungsten carbide and high-speed steel resulted in significant improvements in machining productivity and serve as the predecessors of CVD diamond-coated cutting tools. E.
Commercialization of CVD Diamond Technology for Cutting Tools
The relatively recent development of CVD diamond technology [12–14] has enabled further advances of diamond as a strategic material in the cutting tool industry. Like many new materials, CVD diamond has had its technical and commercial challenges. Today, however, the industry has achieved a track record of success and continues to develop processes that ensure consistency, reliability, and performance [1,2,15–23]. Presently, CVD diamond cutting tools are commercially available in two primary product forms: ‘‘thick-film’’ freestanding CVD diamond cutting tool tips and ‘‘thin-film’’ CVD diamond coatings. 1.
Thick-Film CVD Diamond Cutting Tools
Thick-film CVD diamond is vapor-deposited diamond manufactured in the form of solid, freestanding wafers of pure, binder-less diamond with typical thickness ranging from 150 to over 1000 m. For most cutting tool uses, CVD diamond wafers are laser cut into hundreds of mirror-finish polished CVD diamond tips that are subsequently attached to cutting tools in a manner similar to cobalt-sintered polycrystalline diamond (PCD). The development of successful brazing techniques for thick-film diamond in the early 1990s bypassed the adhesion and substrate selection challenges of thin-film CVD diamond tools [8,23]. Figure 2a contains an image of two large-diameter, as-deposited, freestanding diamond wafers in addition to a collection of mirror-finished freestanding diamond materials. Figure 2b shows examples of finish-ground thick-film CVD diamond tooling (also known as diamond sheet tools [11,23]) adjacent to examples of aluminum workpiece materials. Thick-film CVD diamond cutting tools are engineered to compete
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Figure 2 Thick-film CVD diamond is vapor-deposited diamond that can be manufactured in the form of large-diameter, freestanding wafers of pure, binder-less diamond. The larger wafer shown in the background of image (a) is approximately 175 mm in diameter. For most cutting tool uses, CVD diamond wafers are laser cut into hundreds of mirror-finish polished CVD diamond tips [(a), foreground], which are subsequently attached to cutting tools (b) in a similar manner to cobalt-sintered polycrystalline diamond (PCD). Freestanding CVD diamond cutting tools are used for a wide range of applications including aluminum automotive component machining (b). (From Ref. 1.)
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with conventional PCD tools and are useful in a wide range of applications. As with PCD, thick-film CVD diamond tools are normally made with oversized tips that can be reground multiple times by diamond tool fabricators. Approaches to engineering thick-film CVD diamond-tipped tools are discussed further in Sec. II.D.2. 2.
Thin-Film CVD Diamond-Coated Cutting Tools
Thin-film diamond is CVD diamond that is deposited as a continuous, conformal coating directly onto a suitable substrate. The ability to deposit polycrystalline diamond coatings directly onto cutting tools enables the properties of diamond to be captured in shapes previously not available from nature or the diamond industry. In addition, the complete coating can eliminate the depth of cut limitations of thick-film CVD tools and PCD-tipped tools. Commercial diamond coatings generally fall in the thickness range of 5 to 35 m. The diamond-coated tungsten carbide end mill depicted in Fig. 3 shows the complex shape, high-aspect-ratio coating capability presently available from the CVD diamond industry. Approaches to engineering thin-film CVD diamond-coated tools are discussed in Sec. II.D.1. Thin-film adhesion issues, challenges, and solutions are the topic of Sec. III.B.
Figure 3 Thin-film CVD diamond is vapor-deposited diamond deposited directly onto a component such as a cemented carbide cutting tool. This image contains a diamond-coated tungsten carbide rotary tool showing the complex shape coating capability now available in the CVD diamond industry. The end mill shown is 20 mm in diameter and 160 mm in overall length with a coated flute length of 110 mm. (From Ref. 15.)
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ISSUES AND CHALLENGES IN CVD DIAMOND CUTTING TOOL DESIGN AND MANUFACTURE
Capturing the unique properties of diamond in the design of CVD diamond cutting tools requires a knowledge of the application range and a fundamental understanding of the processes available to manufacture the tool. A broad-scope assessment of the complex properties-process-application interaction is critical because the performance advantage of a CVD diamond tool —if any—will depend not only upon the properties of the diamond itself but also on whether the chosen tool design and manufacturing process effectively capture the desired properties for the application of interest. A.
Critical Properties of CVD Diamond for Mechanical Applications
The key properties of CVD diamond for mechanical applications are not limited to the often-referenced hardness of diamond. The true challenge in CVD diamond tool manufacturing is capturing this potential within a functional tool that delivers reproducible performance. Table 1 provides a comparison of selected mechanical and thermal properties for PCD, single-crystal diamond, cemented carbide, and silicon nitride. Selected properties are summarized further in the remainder of this section. 1.
Hardness and Young’s Modulus
Diamond’s unmatched hardness and elastic modulus contribute to excellent wear resistance. Unlike metal-bonded PCD compacts and cemented carbides, diamond-like carbon (DLC), and many other hard coatings, CVD diamond maintains its hardness and wear resistance at elevated cutting temperatures. Diamond’s hot hardness is due in part to the high-temperature structural stability of its dense, polycrystalline construction coupled with the lack of metal binders that soften at elevated temperatures. 2.
Coefficient of Friction
CVD diamond’s low friction coefficient contributes to lower cutting forces, lower spindle power consumption, and less frictional heating in many cutting applications. The reduced cutting temperature that results from diamond’s high lubricity is viewed as a major advantage in applications where coolant elimination and higher speed operation are being considered. This property is not included in Table 1 because measured values of the coefficient of friction vary significantly depending on environment, surface roughness, and measurement technique. The coefficient of friction of polished CVD dia-
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Comparison of Thermal and Mechanical Properties for CVD Diamond and Other Hard Materials
Property Indentation hardness (GPa)
Young’s modulus (GPa) Tensile strength (MPa)
CVD diamond
Natural diamond
85–100 Technique unspecified Commercial data Ref. 90
50–100 Technique unspecified Commercial data Ref. 90
78.5 ⫾ 18.6 Vickers indentation Commercial data Ref. 91
56–102 Knoop indentation Normal load of 1–500 kg Ref. 92 1000–1100 Commercial data Ref. 90 1050–3000 Orientation dependent Measurement technique unspecified Typical commercial data Ref. 90
1000–1100 Commercial data Ref. 91 Growth face in tension: 500–800 (0.5 mm thick) Three-point bend geometry Estimated from Fig. 11 of Ref. 21 Nucleation face in tension: 800–1300 (0.5 mm thick) Three-point bend geometry Estimated from Fig. 12 of Ref. 21
PCD 50 Technique unspecified Commercial data Ref. 90
WC-Co cemented carbide 7.8–23.5 Vickers indentation Typical range Ref. 93
Si3N4 ceramic 7.8–17.7 Vickers indentation Range based on commercial data from two suppliers
21.5 Knoop indentation Ref. 92
776 Commercial data Ref. 90 1260 Measurement technique unspecified Commercial data Ref. 90
400–700 Ref. 93 700–1500 Typical range Ref. 93
120–330 Typical data Ref. 94 586–843 Range observed for three commercial grades Four-point bend with crosshead speed of 0.004 cm/s NIST evaluated data Ref. 95
Compressive strength (MPa)
16,000 Commercial data Ref. 90
7600 Commercial data Ref. 90
3000–7000 Typical range Ref. 93
>4000 Commercial data ASTM C773 Ref. 97
1.3 Commercial data Ref. 90
8680 average 16,530 maximum Natural octahedron containing no visible flaws or inclusions 10⫺6 m2 estimated load area Ref. 96 2.9 Commercial data Ref. 90
Transverse rupture strength (GPa)
1.2 Commercial data Ref. 90
0.6–3.0 Typical range Ref. 93
5.3–7.0 Commercial data Ref. 91 500–2200 Commercial data Ref. 90
3.4 Commercial data Ref. 90 600–2200 Commercial data Ref. 90
8.81 Commercial data Ref. 90 560 Commercial data Ref. 90
500–1100 Commercial data Ref. 90
600–1100 Commercial data Ref. 90
200 Commercial data Ref. 90
5–30 Typical range Ref. 93 20–120 Typical range for manufacturers quoted values Ref. 93 —
0.76–0.85 Commercial data Ref. 4, p. D116, Kennametal, Inc. 1.5–8.5 Typical values Ref. 94 7–43 Typical data Combined range for four material types Ref. 94 —
Fracture toughness [MPa(m)1/2] Thermal conductivity at room temperature (W/m ⭈ K) Thermal conductivity at 200⬚C (W/m ⭈ K)
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Property Thermal expansion coefficient (ppm/K in 100–250⬚C range) Thermal expansion coefficient (ppm/K at 500⬚C) Thermal expansion coefficient (ppm/K at 1000⬚C) Mean thermal expansion coefficient (ppm/K in denoted range)
CVD diamond
Natural diamond
PCD
WC-Co cemented carbide
Si3N4 ceramic
1.2 Commercial data Ref. 90
1.2 Commercial data Ref. 90
4.2 Commercial data Ref. 90
—
—
3.8 Commercial data Ref. 90
3.8 Commercial data Ref. 90
—
—
—
4.5 Commercial data Ref. 90
4.5 Commercial data Ref. 90
6.3 Commercial data Ref. 90
—
—
—
—
—
4.5–8.5 Typical range for 0–800⬚C temperature range Ref. 93
2.8–3.7 Typical data for 20–1000⬚C range Ref. 94
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mond is generally superior to that of cemented carbides [24] and advanced ceramics [25] and is similar to values attained for single-crystal diamond (e.g., 0.05 [24]). The high lubricity is often compared with that of fluoropolymers.* 3.
Thermal Conductivity
High-quality diamond has the highest thermal conductivity of any material. As Sec. III.A will explain, the use of the highest crystalline quality CVD diamond is not necessarily the best overall solution for mechanical applications of diamond. Regardless, it has been demonstrated that freestanding diamond that has been engineered for mechanical applications can offer thermal conductivity over five times that of tungsten carbide and over 50% greater than that of PCD.† Although diamond’s high lubricity helps to avoid frictional heating effects, CVD diamond’s excellent thermal conductivity relative to other cutting tool materials acts as an additional safeguard in maintaining lower cutting temperatures by dissipating heat generated in the cutting zone. 4.
Thermal Expansion Coefficient
CVD diamond and single-crystal diamond have essentially identical thermal expansion characteristics, as Table 1 delineates. Diamond has a very low thermal expansion coefficient relative to most other cutting tool materials such as cemented carbide and particularly high-speed steel, especially in the 100 to 250⬚C range. This property can be beneficial in controlling the positional accuracy of a cutting tool relative to a workpiece during precision machining. Because of the high-temperature nature of commercial CVD diamond coating and brazing processes, attachment of CVD diamond to cutting tool materials such as cemented carbide requires strict attention to the thermal expansion mismatch. The residual thermal stresses that result from high-temperature attachment necessitate good interfacial adhesion for both braze fabrication of thick-film CVD diamond and direct thin-film coating. Section III.B contains a detailed discussion of diamond adhesion issues for thin-film cutting tools.
*According to E.I. du Pont de Nemours and Company, the coefficient of friction of Teflon is generally in the range of 0.05 to 0.20, depending on the load, sliding speed, and particular Teflon coating used. †Comment is based on Norton Diamond Film’s ‘‘PolyTurn’’ and ‘‘DiamaPak’’ cutting tool products, which were retracted from the market upon business closure.
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Fracture Strength
As Table 1 depicts, the compressive strength of CVD diamond is generally superior to that of single-crystal diamond and PCD. Unfortunately, the pure polycrystalline structure results in a material that is generally inferior in tensile strength. Even so, the tensile strength of freestanding CVD diamond is sufficient for a wide range of industrial applications, especially when incorporated into tool designs that accommodate the unavoidable fact that diamond is a brittle, ceramic-like material. The measured values of the Weibull modulus [21] indicate that the breadth of the strength distribution for bulk CVD diamond is similar to that for advanced ceramics such as silicon nitride. The observed values ensure that CVD diamond can yield the reliable, reproducible tensile strength necessary in many mechanical applications. 6.
Fracture Toughness
Fracture toughness is a property of a material that is frequently defined as resistance to crack propagation. High fracture resistance is necessary for a cutting tool material to survive the impact and fatigue that can result from aggressive cutting operations such as milling or interrupted cutting or when machining workpiece materials containing material inhomogeneities such as hard phases, fibers, or particles. As shown in Table 1, the fracture toughness values of bulk, freestanding CVD diamond are comparable to those of advanced ceramic materials such as silicon nitride. Furthermore, because the microstructure of polycrystalline CVD diamond can provide a crack deflection mechanism [26–28], proper microstructural engineering can yield fracture toughness values exceeding those of single-crystal diamond and approaching values common to cobalt-sintered PCD. Section III.A includes approaches to optimizing the bulk properties of CVD diamond including fracture toughness. 7.
High-Temperature Phase Stability
Black, graphite-like coatings have been observed on diamond surfaces exposed to atomic oxygen at temperatures as low as 625⬚C [24]. However, diamond oxidation occurs at very low rates in air below 800⬚C and generally does not limit the application range of CVD diamond tooling. CVD diamond is thermally stable to temperatures of at least 1200⬚C in vacuum. By contrast, PCD is particularly susceptible to thermal degradation at temperatures above 700⬚C [29]. The cobalt binder phase can catalyze graphite formation at the sintered contacts between the diamond grains. This high-temperature phase instability, which can be induced by excessive cutting temperatures or im-
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proper brazing or grinding procedures, is known to result in a general degradation of the mechanical properties of the PCD cutting tool tip. 8.
Chemical Corrosion Resistance
The chemical reactivity of carbon limits the use of diamond tooling in ferrous metal and superalloy machining. All forms of diamond tooling are at risk of corrosion or graphitization induced by machining iron-, nickel-, and cobalt-containing metals at elevated cutting temperatures [30]. Molten manganese, chromium, and the platinum group of metals are also true solvents for all forms of carbon including diamond [24]. Other metals capable of reacting with diamond include the carbide-forming transition metals titanium, tantalum, tungsten, and zirconium. Even so, diamond’s superlative combination of properties still maintains it as the preferred form of tooling for selected applications such as bimetallic machining of aluminum–cast iron deck faces for automotive engines [5]. With the exception of the metallic interactions just noted and hightemperature oxidizing environments, pure forms of diamond are extremely chemically inert. Unfortunately, the corrosion resistance of CVD diamond is a property that is often overlooked. For example, the metallic constituents of PCD and cemented carbides limit tool lifetime in corrosive polymer machining. Although the corrosion resistance of ceramic cutting tools can offer measurable improvements over traditional tooling, the advent of CVD diamond technology now allows the production of tooling with complex geometry that can offer both corrosion resistance and unmatched abrasive wear resistance. B.
The Critical Combination of Abrasive Wear Resistance and Fracture Toughness
In order to meet new machining requirements for difficult-to-machine materials, improve productivity, and hold high workpiece tolerances, new cutting tool materials must have both suitable fracture toughness and high abrasive wear resistance. The importance of fracture toughness, a measure of the resistance to crack propagation, was discussed in Sec. II.A.6. However, the specific abrasive wear resistance of a CVD diamond tool is very application specific and, therefore, was not discussed in Sec. II.A. Abrasive wear resistance is the result of a complex interaction between the properties and characteristics of the diamond tool, the workpiece material being machined, and the nature of the specific application. Improved abrasive wear resistance allows machinists to maintain tight tolerances with minimal tool or machine adjustments and reduces machine downtime associated with rapid tool wear and frequent tool changes.
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Figure 4 depicts schematically the trade-off between fracture toughness and abrasive wear resistance in cutting tool materials. For example, advanced ceramic cutting tool materials, known to be very hard and resistant to abrasive wear, tend to be lower in fracture toughness. Metallic cutting tools such as high-speed steel, known to be high in fracture toughness, tend to have relatively low resistance to abrasive wear. Although cemented carbide does not have the superior toughness of high-speed steel, it combines an abrasive wear resistance approaching that of ceramic materials and fracture toughness that is sufficient for a wide range of heavy-duty industrial machining needs. In most applications, the exceptional fracture toughness of high-speed steel is not necessary or worth the inherent trade-off in abrasive wear resistance. Figure 4 also shows that ceramic-like CVD and PVD coatings can markedly enhance the abrasive wear resistance of both tungsten carbide and high-speed steel tooling while the critical fracture toughness of the underlying tool substrate is maintained. From the success of hard coatings in the
Figure 4 Abrasive wear resistance and fracture toughness trade-offs for common cutting tool materials.
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cutting tool market, it is clear that the added hardness and wear resistance of diamond-coated tungsten carbide offer a well-balanced fracture toughness-wear resistance combination. The unique combination expected for CVD diamond-coated cutting tools is depicted at the top center of the figure. Because fracture toughness is a bulk mechanical property and the underlying substrate of a diamond-coated carbide tool is tougher than bulk CVD diamond, the diamond-coated carbide concept offers the potential to exceed the fracture toughness of bulk, freestanding CVD diamond. However, as discussed in Sec. III.B of this chapter, the composite properties of a diamondcoated carbide are captured in the tool design only if the adhesion strength exceeds the forces of the application. The abrasive wear resistance of the cobalt-sintered diamond in PCD can offer as much as 100 times greater tool life over cemented carbide [5]. By contrast, as Fig. 4 shows, PCD-tipped cutting tools are usually inferior in fracture toughness to cemented carbides. Fortunately, the fracture toughness of PCD is sufficient to avoid cutting edge chipping in most rigid machining situations. In fact, as the operating speeds and feeds move beyond the material limits of traditional cutting tool materials in aggressive highproductivity operations, the long life and predictable finish advantages of diamond tooling are often amplified.
C.
Application Characteristics That Affect CVD Diamond Cutting Tool Performance
Every cutting tool application is unique, and industrial cutting tools need to be designed for the worst-case scenario within the targeted range of applications. There are many factors that can directly or indirectly affect cutting tool performance. The feasible application range of a CVD diamond cutting tool design may not be clear until a tool is thoroughly tested in the field. As with any cutting tool, success in a specific operation on a specific machine tool does not ensure the same results with the same tool on the same type of machine at a different location. Because of the complexity of machining processes, this section is provided as an overview of the application characteristics that should be factored into the design and manufacturing processes used to produce CVD diamond cutting tools. 1.
Impact of Machine, Fixture, and Operating Conditions on Performance
Factors such as spindle balancing and rigidity, workpiece vibration, or even poor tool handling practices can dictate whether or not a tool will offer
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a performance advantage in a specific application. Clearly, the type of machining operation (e.g., turning, boring, milling, drilling) and machine operating parameters (e.g., cutting speed, cutting feed rate, depth of cut) have a major influence on the nature of the forces on the CVD diamond cutting edge and the diamond-tool interface. For instance, machining operations where the cutting edge is not in constant contact with the workpiece (e.g., interrupted turning, milling, or end milling) generally require more robust design than continuous cutting operations (e.g., turning, boring, drilling). The impact-like forces generated in interrupted operations require good material strength and fracture toughness to minimize the damage of a CVD diamond tip or diamond coating during use. Other important application details include the coolant type (e.g., dry, mist, flood), chip-forming behavior, and efficiency of chip evacuation from the cutting area. Because CVD diamond (like other forms of diamond) is capable of operating at high cutting speeds and feed rates, machine tools with limited ranges of speed and feed can significantly limit the performance potential of diamond. By contrast, the wear resistance of many traditional cutting tool materials can degrade rapidly when machining processes are put into ‘‘fast forward’’ at high speeds and feeds. As a demonstration, Cline [31] and Tanikella et al. [32] have reported the results of end mill machining tests that were performed on SiC-containing ‘‘bisque ceramic’’ that was heat treated (‘‘presintered’’) at temperatures greater than 1200⬚C. The workpiece material and machining conditions were selected in an effort to accelerate the diamond wear mechanism common to abrasive applications such as ‘‘green ceramic’’ (unfired polymer-ceramic compacts) and graphite machining. The results of the machining tests showed the dramatic performance advantage of thin-film CVD diamond-coated end mills at the selected machining conditions. Figure 5a and b show that the uncoated tungsten carbide and TiNcoated tungsten carbide tools failed within 16 cm of machining (two 7.6cm-long passes). The failure was highlighted by excessive abrasive wear, which was enhanced by frictional heating of the tools. The PVD TiN coating offered no clear advantages under these machining conditions. Optical pyrometry of the carbide and TiN-coated carbide tools indicated temperatures exceeding 900⬚C. By contrast, a diamond tool that had already run 80 previous passes was barely warm to the touch after eight sequential passes at the same machining conditions. Freeze-frame images of the 88th pass are shown in Fig. 5c. Both the TiN-coated and uncoated carbide tooling showed a strong sensitivity to density variations in the presintered ceramic workpiece. In the case of one uncoated carbide tool, the excessive wear and overheating reduced the 9.5 mm cutting diameter of a new tool by over 2.5 mm within less than 23 cm of machining. More details of this case study are summarized in Sec. IV.B.2.c.
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Figure 5 Effect of high cutting speeds on various cutting tool materials: sequential freeze-frame images showing presintered ceramic machining using (a) an uncoated tungsten carbide end mill, (b) a TiN-coated carbide end mill, and (c) a CVD diamond-coated end mill. The uncoated and TiN-coated tungsten carbide tools failed within 16 cm of machining due to excessive wear and frictional heating. The diamond-coated tool operated at cutting temperatures at least 500–1000⬚C lower than the uncoated equivalent and was still worthy of continued testing after more than 670 cm of machining. (From Ref. 31.)
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Figure 5
2.
Continued
Impact of Workpiece Material Characteristics on Performance
The relative performance advantage of one tool over another is dependent on the workpiece material and the requirements for finished part quality. The characteristics of the workpiece such as abrasiveness, corrosiveness, and microstructural homogeneity can significantly affect the utility of a diamond tool in machining operations. Although less aggressive applications such as graphite machining will not always differentiate between tools of varying integrity, applications that generate higher, impact-like cutting forces can rapidly differentiate between tools that offer predictable abrasive coating wear and tools that fail unpredictably due to film delamination. The characteristics of diamond tooling are of particular interest in machining composite materials. The unique advantages offered by composites are generally the result of combining two or more dissimilar metallic and/ or nonmetallic materials in an effort to yield combined properties that exceed those of the individual constituents. For instance, glass fiber–reinforced plastics (GFRPs) combine the stiffness of the reinforcing fibers with the lowcost, lightweight characteristics of an otherwise pliable polymer matrix. Other useful composites, such as high-silicon aluminum alloys, metal matrix
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composites (MMCs), and structural aerospace composites, offer advantages ranging from weight reduction to wear resistance and fatigue resistance. By design, the most unique and enabling composites tend to be composed of materials that have diverse properties and, therefore, vastly different machining behavior. The combined results are workpiece materials that can be abrasive, corrosive, and generally difficult to machine with conventional cutting tool materials [2,3,23,33,34]. Table 2 lists some commonly used metallic and nonmetallic composite materials along with their associated machining challenges. Many early efforts to fabricate components from these and other composites using conventional tooling proved to be cost prohibitive due in part to high machining costs. As the industrial demand for innovative composite materials continues to increase, so does the interrelated demand for cutting tools that can endure the challenges of machining these progressive materials. Many of the case studies discussed in Sec. IV are related to this increasingly important class of materials.
D.
Design and Manufacturing Options and Constraints for CVD Diamond
The design and manufacture of diamond-coated cutting tools varies significantly from thick-film, brazed-on CVD-diamond cutting tools. The last step in manufacturing a thin-film, diamond-coated tool is typically the CVD coating process, whereas freestanding CVD diamond wafer fabrication is the first of many steps in the fabrication of a thick-film, brazed-on CVD diamond tool. Many of the process-related advantages of the thin-film CVD diamond coating process are similar to advantages observed with other hard coatings used for cutting tools (see Sec. I.D). On the other hand, the processrelated advantages of the fabrication process for thick-film CVD diamond are more typical of PCD diamond tooling (see Sec. I.C.2). 1.
Approaches to Engineering Thin-Film CVD Diamond-Coated Tools
Today, CVD diamond processes enable direct coating of high-aspect-ratio, complex shapes ranging from specially molded, chip-breaking inserts to the end mill shown in Fig. 3. However, the manufacture of thin-film, CVD diamond-coated cutting tools does not come without unique technical challenges. Most noteworthy are the challenges associated with adhesion of diamond deposits to tungsten carbide—the preferred substrate for diamond
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Examples of Difficult-to-Machine Composite Materials Example
Examples of use
Machining issues
Related CVD diamond performance case studies
Composite: A390 (18% Si in Al) Matrix: Aluminum Reinforcement: Silicon particles Composite: Duralcan (30% SiC in Al) Matrix: Aluminum Reinforcement: Silicon carbide ceramic particles Composite: 25% cobalt tungsten carbide Matrix: Cobalt-based alloy Reinforcement: sintered tungsten carbide grains Composite: Carbonepoxy Matrix: Epoxy polymer Reinforcement: Highdensity carbon fibers Composite: G10 Matrix: Polymer Reinforcement: Glass fibers
Reduced-weight, wear resistant, temperatureresistant pistons
Although aluminum is generally free-machining, hard silicon particles are extremely abrasive
Thick-film inserts: IV.A.1.a Thick-film rotary tools: IV.A.2 Thin-film inserts: IV.B.1.a, IV.B.1.b
Brake rotors Lightweight structures
Although aluminum is generally free-machining, hard SiC ceramic particles are extremely abrasive
Thick-film inserts: IV.A.1.c Thin-film rotary tools: IV.B.2.c
High-fracture-toughness wear parts
Tungsten carbide sintering results in strong bonding or contiguity between WC grains. The WC grains are very hard. At elevated temperatures, the metal binder can react with the carbon in diamond. Carbon fibers are extremely abrasive; grinding is often preferred over cutting due to excessive tool wear
Thick-film inserts: IV.A.1.d
Glass fibers induce abrasive tool wear; polymer can cause corrosive tool wear
Thin-film rotary tools: IV.B.2.c
Material type Hypereutectic or ‘‘high-silicon’’ aluminum
Metal matrix composites (MMCs)
Cemented tungsten carbide
Fiberglassreinforced polymers (FRPs)
Source: Modified from Ref. 2.
Stiff, lightweight support structures for commercial aircraft Strong, lightweight sporting goods Lightweight, insulative circuit boards Low-cost structural composites
Thin-film rotary tools: IV.B.2.c
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Structural ‘‘aerospace composites’’
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coating commercialization. Many early efforts to introduce diamond-coated carbide cutting tools into the aluminum machining market failed or were significantly delayed because of lack of adhesion treatment technologies that could reproducibly yield WC-Co surfaces worthy of CVD diamond processing and subsequent use. Many early development efforts in the CVD diamond cutting tool industry involved the use of advanced ceramic materials as interim substrate solutions because the chemical compatibility and thermal expansion mismatch were more favorable for diamond coating. A variety of substrate materials including SiAlON, binder-less tungsten carbide and, more commonly, silicon nitride and silicon carbide were employed [23]. Although diamond-coated ceramic tools do have utility in selected applications, the inferior fracture toughness of diamond-coated ceramic inserts and end mills generally limited their use to nonmetallic machining operations such as those involving graphite, polymers, green ceramics, and nonmetallic composite materials. a. Interfacial Integrity. Figure 6 contrasts the failure modes of a ‘‘bad’’ tool (Fig. 6a) that has failed by unpredictable, catastrophic film delamination related to poor interfacial adhesion and a ‘‘good’’ tool (Fig. 6b) that has failed by predictable abrasive wear. The relative utility of these two tools is clear. If the diamond adhesion is not suitable for the application, the desired coating properties simply are not captured in the tool design. As discussed in Sec. III.B in detail, proper diamond-coated WC-Co cutting tool design requires strict attention to both the coating properties and the substrate surface properties. b. Morphology and Surface Roughness Characteristics. Today, most thin-film CVD diamond cutting tools are made from well-faceted CVD diamond deposits because of the higher abrasive wear resistance typically offered by coarse-grained diamond. However, relatively smooth, more lubricious, microcrystalline diamond coatings are commercially available for applications where workpiece surface finish is more critical. Most diamondcoated tools are presently marketed in an as-deposited state with no postdeposition finishing. Depending on the interfacial roughness introduced by adhesion treatment, coating thickness, and the morphological features of the coating, the surface roughness and edge blunting of some diamond-coated tools can limit their use to intermediate finishing and roughing applications. The morphological features of CVD diamond are discussed in more detail in Sec. III.A. c. Postdeposition Finishing Options. Postdeposition finishing techniques can be applied to modify the generally rough, matte finish of diamond-coated products. Diamond-coated ceramic tools with mirror-finished
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Figure 6 Contrasting failure modes of CVD diamond coatings: SEM images contrasting the failure mode of two different diamond-coated tool designs after aluminum metalcutting tests. (a) A tool that failed catastrophically due to insufficient coating adhesion. A metallic buildup of aluminum is evident in the delaminated zone on the tool tip where the adhesion-treated carbide was exposed. (b) A tool that offered a suitable level of adhesion for the application and therefore maintained the slow, predictable wear characteristics desired in a diamond-coated component. As reviewed in Sec. III.B, the diamond/substrate interface characteristics have a significant influence on the nature of tool failure. (From Norton Diamond Film/Saint Gobain Industrial Ceramics, Inc.)
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rake faces similar to PCD tools have been produced in the industry.* Oles et al. [23] have reported the success of a CVD diamond buffing technique in reducing insert surface roughness for the flank face of a diamond-coated carbide insert. Turning tests in A390 (18% Si-Al) showed that the buffed flank surfaces substantially improved the workpiece surface finish. However, the turned component finish of the buffed tools was still inferior to the finish generated by finish-ground PCD tools tested in the same study. Kanda et al. [17] have reported progress in grinding rotary tools to enhance cutting edge characteristics and resulting workpiece surface finish. Although the general tendency of the industry is toward smoother surfaces that mimic the finish ground characteristics of PCD and cemented carbide tools, Oles et al. [23] have also reported a ‘‘microscopic chip-breaking’’ advantage of diamond-coated cemented carbide inserts. Milling and turning studies of low-silicon and high-silicon alloys showed the advantages of diamond facets on the chip-forming rake surface in creating a ‘‘highly stressed chip that is more readily broken.’’ Metallic chips in the shape of small ‘‘sixes’’ and ‘‘nines’’ were reported to be more desirable than continuous stringers or coils that are common to smooth rake-face tools such as PCD. As CVD diamond technology advances, it is likely that the use of polishing, grinding, or other finishing techniques for CVD diamond-coated tooling will become more prevalent. Increases in CVD diamond cutting tool demand will provide justification for tailoring film properties such as thickness and surface roughness to specific applications. Such justification is critical to the presence of CVD diamond in the market because it will undoubtedly lead to an expansion in application range for diamond-coated products. 2.
Approaches to Engineering Thick-Film CVD Diamond Tools
a. Design and Manufacture of Thick-Film Tips. The primary technical challenge in the manufacture of freestanding thick-film CVD diamond cutting tool tips is attaining a throughput of freestanding diamond material having the dimensional and mechanical characteristics needed for downstream processes and ultimate use [35]. This is commonly approached through the manufacture of large-area, crack-free, CVD diamond discs of the necessary thickness with minimal thickness variation and minimal bow. Attaining CVD diamond wafers of this quality is exceptionally difficult due to the unique mechanical properties of diamond and the tendency for these *Mirror finish observed on diamond-coated silicon carbide cutting tools once produced by DeBeers.
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properties to vary measurably with changes in deposition conditions such as temperature and gas chemistry. Clearly, factors such as CVD diamond reactor type, the associated diamond deposition area, deposition rate, and vacuum process control capabilities are critical to success in freestanding diamond production. Although this issue has been discussed in various forums [35], many of the details of the industrial methods used to generate large freestanding diamond discs are presently treated as trade secrets and will not be discussed in this text. Regardless of the challenges, pure CVD diamond wafers for mechanical applications have been commercially manufactured in the thickness range of 150 to over 1000 m with diameters up to 175 mm (see Fig. 2a).* By comparison, commercial high-temperature, highpressure PCD manufacture is presently limited to disc sizes of about 74 mm. Cutting CVD diamond components from freestanding CVD diamond wafers is presently done using laser technology as opposed to the electric discharge machining (EDM) ‘‘wire cutting’’ common to the PCD industry. Pure CVD diamond wafers contain no metallic binders and, therefore, do not afford the conductive properties needed in an EDM workpiece. The general acceptance of EDM wire techniques for diamond tool fabrication coupled with the present lack of the necessary laser technology at cutting tool fabricator sites is a significant barrier to commercialization of thick-film diamond tooling. b. Attachment Techniques for CVD Diamond Tips. Most freestanding, thick-film CVD diamond used for cutting tool applications is approximately 0.5-mm-thick polished diamond. The pure diamond can be attached to the tool body using a high-temperature inert vacuum braze (normally >850⬚C). The use of a high-temperature braze ensures good chemical bonding and excellent braze strength. Commercially available forms of freestanding CVD diamond cutting tool tips are generally made from well-faceted deposits that may have crystal sizes exceeding 50 m. The surface roughness of the thick-film, freestanding diamond is usually removed by diamond finishing techniques in an effort to mimic the desirable characteristics of PCD. In most commercial thick-film CVD cutting tool applications, the relatively smooth nucleation surface is used as the cutting surface and the rough, faceted face becomes part of a vacuum-brazed interface typically bonded to tungsten carbide. A flank-face view of a finish-ground CVD diamond cutting tool tip is shown in Fig. 7. Commercially available thickfilm cutting tool tips offer superior wear resistance, which is typical of coarse-grained thin-film diamond deposits, coupled with mirror finish (commonly 1 m), faceted crystallines to fine-grained (diameter 1 m) faceted crystallites to fine-grained (diameter Figure 11 Comparison of intergranular and intragranular fracture in CVD diamond. SEM micrographs of the fractured cross section of thin, CVD diamond films grown on WC-Co substrates show the dramatic variations in fracture characteristics that can result from simple changes in CVD diamond process conditions. A highly oriented film seen as the dark layer in the top of the 2000⫻ image in (a) displays weak grain boundaries and a fracture mode that is intergranular. A randomly oriented film seen as the dark layer making up most of the 4000⫻ image in (b) displays strong grain boundaries and a fracture mode that is nearly 100% intragranular. (From Ref. 74.)
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processes operate at high temperatures, diamond coating requires special attention to substrate chemistry and general process control that is not as critical for most other coatings. Diamond coating challenges are amplified by the fact that the primary market demand is for diamond coatings on existing grades of tungsten carbide that contain not only cobalt but also various grain growth inhibitors and metallic impurities. The high-temperature reactivity of diamond with the cobalt-based binder phase of tungsten carbide cutting tools generally prevents the direct, high-temperature deposition of CVD diamond directly onto ‘‘off-the-shelf’’ tungsten carbide grades. Even after high-temperature chemical compatibility of the interface is ensured using adhesion treatment processes, the thermal expansion mismatch between the film and substrate often requires attention as it can be the source of very large residual stresses. Ignoring any one of these factors can lead to poor interfacial integrity and may result in delamination of the film or coating during use. Because overcoming diamond adhesion challenges is part of the critical path to successfully manufacturing a thin-film CVD diamond cutting tool, this section is devoted specifically to adhesion issues. Section III.B.1 is an overview of the specific adhesion challenges in coating cutting tools with a focus on commercial grades of tungsten carbide. Section III.B.2 includes various approaches that have been used to overcome adhesion barriers in an effort to capture the enabling characteristics of diamond-coated tungsten carbide. Section III.B.3 contains a case study on the development of a specific adhesion treatment intended for the manufacture of diamond-coated metalcutting inserts. 1.
Diamond-Substrate Characteristics Necessary for Adhesion
Adhesion strength may be loosely defined as the force or energy required to separate two objects. It may be thought of as the energy necessary to break the bonds at an interface, thus driving the extension of an interfacial crack. Extension of such a crack is referred to as interfacial crack propagation. It is interfacial crack propagation that leads to failure of an interface and ultimately yields visually apparent signs of component failure referred to as ‘‘delamination,’’ ‘‘spalling,’’ or ‘‘flaking.’’ Using this definition, interfacial failure is defined not solely by a fracture resistance parameter (‘‘interfacial fracture toughness’’) uniquely related to the bonding across the interface but by the ‘‘adhesion strength,’’ which is determined by the combined influences of interfacial fracture toughness, strength controlling defects, and residual stresses [52]. The interfacial fracture toughness is related to the type and density of interfacial defects (cracks, voids, weak interfacial
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phases, etc.), the density of chemical bonds between the diamond film and substrate at the interface, and the presence of an interfacial crack deflection mechanism [52–54]. Therefore, maximizing adhesion strength for CVD diamond-coated cutting tools requires minimizing specific interfacial defects such as voids, cracks, or weak interfacial phases; maximizing chemical bond density; and providing a crack deflection mechanism that increases the energy necessary to drive an interfacial crack. In order to minimize the interfacial defects and maximize the chemical bond density in diamond-coated WC-Co, it is essential to maximize diamond nucleation density, ensure that the interface is chemically stable, and suppress the formation of ‘‘soft’’ nondiamond phases at the interface. As will be discussed in Sec. III.B.1.d, toughness-enhancing crack deflection characteristics can be engineered into the diamond/WC-Co interface. It should be emphasized that, in practical applications, the ‘‘adhesion strength’’ (in quotation marks to emphasize the potential difference in failure mechanism) is a measure of the utility of the tool in the application and its resistance to decohesion during use. That is, a quantitative assessment of adhesion strength is meaningless unless the conditions mimic the specific failure mode observed in the application. Using this alternative definition, the failure mode commonly described as delamination, spalling, or flaking is failure that is located at or near the interface. This distinction encompasses crack propagation within the nucleation layer of the diamond film as well as failure within the free-surface grains of the substrate such as the adhesion-treated region of a diamond-coated WC-Co component. Although this definition is of much more practical utility, it necessitates the use of semiquantitative failure analysis to verify the specific mechanism and location of the ‘‘interfacial’’ failure. a. Nucleation Density of Diamond. Effective nucleation of CVD diamond onto foreign substrate materials is notably more difficult than nucleation of other hard coatings produced by CVD or PVD processes. Unlike other hard coating processes, ‘‘diamond seeding’’ processes are often employed to enhance the nucleation density of diamond [55]. Heteroepitaxially grown diamond films begin as submicrometer-sized nuclei formed on the top of a nondiamond substrate. The physical characteristics of the nuclei are dependent on the deposition conditions and the composition of the underlying substrate. As schematically shown in Fig. 12, subsequent growth normally occurs by addition of atoms to already formed nuclei rather than by continued formation of new nuclei. If the individual nuclei are separated by a ‘‘large distance,’’ interfacial voids may form when the individual nuclei coalesce and form a continuous layer. By maximizing the nucleation density, the inclusion of internucleic voids at the interface may be minimized. Be-
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Figure 12 CVD diamond nucleation and growth. On many nondiamond surfaces, growth of CVD diamond films occurs by a Volmer-Weber or ‘‘island growth’’ type of nucleation/growth mechanism. In this process, a finite number of nuclei are formed during a nucleation stage (a). Following this stage, adatoms are incorporated into the individual nucleation sites until the islands have grown to some maximum diameter at which their boundaries intersect (b). If the density of diamond nuclei is low and individual nuclei are separated by large distances relative to their size, interfacial voids will form when the nucleation sites have grown to a size at which their diameters coalesce (b). The size of these interfacial or internucleic voids may be minimized (assuming the same nucleation and growth mechanism) by increasing the nucleation density as shown in (c). (From Ref. 74.)
cause these voids may act as crack nucleation sites, they can lower the critical stress necessary for crack propagation at the interface. Increasing the nucleation density also reduces the time necessary for the nuclei to coalesce into a continuous film. Many methods of optimizing the nucleation density for a range of substrate materials have been presented in the CVD diamond synthesis technical literature [55–60] and the details are not discussed in this chapter. b. Interfacial Chemical Bonding. The chemical composition of some substrate materials can impair or negate the ability to form strong bonds between the film and substrate, especially with the high-temperature nature of CVD diamond processes. Physical or mechanical adhesion alone is simply not sufficient to ensure performance. Chemical bonding is necessary to counter the interfacial stresses induced by the thermal expansion mismatch and the applied load of cutting tool operations. Several approaches
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may be taken to optimize the chemical bonding between the film and substrate. In general, the most effective means of maximizing bond density is by minimizing the mismatch between coefficients of thermal expansion and lattice spacings between the film and substrate. A more detailed consideration of this topic may be found in Allen [61]. For the specific case of CVD diamond films on carbide-forming substrates (Ta, Cr, Ti, W, etc.), a carbide ‘‘interphase’’ may be formed that can relax the lattice mismatch between the film and substrate [62,63]. However, in the case in which amorphous carbon or other nondiamond material is used to enhance the nucleation density, these nondiamond layers may form a weak interfacial layer suppressing the strength of chemical diamond-substrate bonds. Therefore, deposition process conditions must be carefully controlled in order to suppress the formation of a weak interfacial nondiamond phase or, if this phase cannot be avoided, to minimize its thickness. c. Substrate Chemical Composition. Because cobalt catalyzes the formation of nondiamond carbon (sp, sp2-bonded carbon) in the diamondCVD growth environment, it is important that the binder phase be absent from the substrate surface during nucleation [64,65] and that diffusion of the binder phase from the bulk to the free surface is suppressed [66]. It is believed that the presence of Co at the growth surface or at a region of already formed diamond can perturb the phase stability of the deposit. Under typical CVD growth conditions, this may result in a phase transformation from diamond to graphite or other nondiamond carbon phases. The solubility of carbon in cobalt under the elevated temperatures of CVD diamond deposition conditions may drive the dissolution of the carbon at the interface into the binder phase [67]. For this reason, the binder phase composition and the proximity of the binder phase to the diamond film may also affect the phase stability of the interface when subsequently exposed to the hightemperature diamond deposition conditions. d. Interfacial Microstructure. Consider a film-substrate system as illustrated in Fig. 13, which is characterized by an ‘‘atomically smooth’’ interfacial microstructure. Because adhesion strength is related to the density of chemical bonds that are present at the interface, the only resistance to interfacial crack propagation in this case is provided by the strength of these chemical bonds. If we consider the interface to contain residual stress and other weakening characteristics such as internucleic voids, the applied force required to drive an interfacial crack is further reduced. Now consider the film-substrate system as illustrated in Fig. 14, which is characterized by a rough interfacial microstructure. Recalling the previous definition that adhesion strength is a measure of the force or energy required to drive a crack along the interface between the film and substrate, this
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Figure 13 Schematic representation of a smooth substrate-film interface system. Interfacial crack extension or propagation is resisted only by chemical bonding between the film and substrate. If an additional stress is applied to a film that has a preexisting crack at the interface (a), this crack will extend along the interface resulting in delamination (b). The adhesion strength of films that display this type of interface is limited by the strength and density of the chemical bonds, residual stresses, and the type and density of interfacial defects such as cracks or voids. (From Ref. 74.)
energy needed is significantly increased for a system that incorporates a nonporous, ‘‘rough’’ interface. This is due to several factors. Generally, the diamond-WC interface (assuming that no Co or Co interaction is present) can be considered a brittle interface. In brittle materials, crack propagation occurs very rapidly due to high crack velocities. Studies of single-crystal diamond have shown that the crack velocities exceed 5000 m/sec [68]. Fortunately, the same crack deflection phenomenon described for the bulk diamond in Sec. III.A.3 is also applicable to the diamond-substrate interface. For a crack to extend along the interface, it must follow a tortuous path as it is redirected along the rough interface as illustrated in Fig. 14. Because of the high crack velocity, the lowest energy crack path is most likely a linear path into the film or into the substrate. Extension of the crack into the film or into the substrate may pin the crack tip and arrest the interfacial crack, effectively toughening the interface. In addition, because the interfacial surface area between the film and substrate is much greater for rough interfaces, the increased density of chemical bonds further enhances adhesion strength.
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Figure 14 Schematic representation of a rough substrate-film interface system. In addition to the increase in surface area available for chemical bonding, this type of interface may provide a tough, crack deflection mechanism. In brittle interfaces characterized by little plasticity, crack extension is redirected from the interface (a) and may be arrested by the film or substrate (b). Thus, the increase in chemical bond density in addition to the toughening mechanism results in an increase in resistance to interfacial crack propagation and delamination. (From Ref. 74.)
2.
Selected Approaches to Solving the Tungsten Carbide Adhesion Problem
a. Wet-Chemical Etching Processes. Wet-chemical etching techniques have been commonly used as an adhesion treatment approach targeted at removing the metallic binder phase from the surface of WC-Co substrates. For example, Kikuchi et al. [69] proposed the formation of an ‘‘etch layer’’ extending between 0.1 to 1.0 m into a WC-Co substrate with a 1–4 wt% Co content. Etching methods successfully allow the nucleation and growth of diamond films on the etched surface without the preferential deposition of graphite. However, most wet-chemical techniques require stringent process control to ensure utility of the diamond-coated part. The binder phase, which is located at the interstices of the hard, WC grains, provides the toughness characteristic of ductility to the WC-Co material. It also serves to ‘‘cement’’ the hard, somewhat contiguous, WC ‘‘skeletal structure’’ together to form a composite material with excellent mechanical properties. Chemical removal of the binder phase to a depth greater than the dimension of the free-surface WC grains essentially removes the cement and generally results in the formation of an embrittled layer at the surface of the WC-Co
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part. In the presence of residual thermal stresses or applied stresses encountered during use, failure of the interface by WC grain ‘‘decohesion’’ or by crack extension in this embrittled area may result in delamination of the film as illustrated in Fig. 15. Diamond-coated tungsten carbide components that have failed by this mode characteristically have WC grain pullout from the WC-Co substrate onto the back side (interface/nucleation side) of the delaminated diamond film. An example of this is shown in Fig. 16. By contrast, the removal of the binder phase to a depth less than the dimension of the free-surface WC grains may allow too much chemical interaction between the diamond and binder phase at the high temperatures typical of traditional CVD diamond synthesis. This interaction results in dissolution of the diamond and diffusion of the carbon into the metallic binder phase as well as diffusion of the binder phase into the diamond as illustrated in Fig. 17. This effectively reduces the interfacial contact area
Figure 15 Overetched WC-Co surfaces yield WC-Co/diamond interfaces that are susceptible to mechanical failure. (a) A schematic representation of the cross section of a WC-Co substrate. The goal of most wet-chemical etch processes is to remove the binder phase from a region near the free surface of the WC-Co surface. If the etch depth extends further into the substrate than the top layer of the WC grains (b), the toughness of the interface will be severely degraded. Following diamond deposition on this surface (c), the diamond-coated part will cool from deposition temperature. Because of the mismatch in the coefficient of thermal expansion between the film and substrate, the interface will be in tension due to residual thermal stresses. In the presence of further applied stresses, crack extension along the interface between the near-surface WC grains and the unetched region of the substrate (d) can lead to failure of this region. This is sometimes referred to as grain pullout as the interface side of the diamond film contains WC grains that were pulled from the substrate (e). (From Ref. 74.)
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Figure 16 SEM micrograph of the interface side of a diamond coating that delaminated due to the failure mode shown in Fig. 15. The bright contrast, faceted grains were verified to be WC grains using energy dispersive spectroscopy (EDS). Detection of these WC grains on the interface side of the diamond film verifies that failure of the interface region occurred by crack extension in the region of the substrate in which the binder phase was removed. (From Norton Diamond Film/ Saint Gobain Industrial Ceramics, Inc.)
between the film and substrate and sacrifices the mechanical integrity of the diamond-coated component. This failure mode has also been characterized by SEM and energy dispersive spectroscopy (EDS) as presented in Fig. 18. Metal machining operations induce cutting forces that often exceed the level of adhesion offered by etching processes. Extensive cutting tests and interfacial analysis of diamond coatings on pre-etched WC-Co surfaces indicate* that etched interfaces of common WC-Co grades offer an insufficient level of adhesion for reproducible performance in applications such as aluminum machining. The adhesion limitations of etched WC-Co are especially evident under more aggressive machining conditions that induce impact forces on the interface. Notable examples include milling, interrupted turning, and the machining of workpiece materials with structural inhomogeneities (e.g., high-silicon aluminum, MMCs). Other testing has indicated that good etching process control can allow production of diamond films on *Comment is based on studies performed within and monitored by Norton Diamond Film prior to business closure.
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Figure 17 Underetched WC-Co surfaces yield WC-Co/diamond interfaces that are susceptible to chemical degradation during coating. (a) Schematic representation of the cross section of a WC-Co substrate. (b) The goal of most wet-chemical etch processes is to remove the binder phase from a region near the free surface of the WC-Co surface. (c and d) If the etch depth does not extend sufficiently far into the substrate, the binder phase and diamond film may interact under the high temperatures typical of most CVD diamond growth conditions. The Co can induce a phase transformation from diamond to graphite. The carbon then diffuses into the binder phase as the cobalt diffuses into the film. The resulting interfacial porosity results in poor bonding between the film and substate. (e) Depending on the magnitude of the damage, delamination of the film can occur because of residual thermal stresses or the applied stresses of the application. (From Ref. 74.)
> Figure 18 Structural and chemical analysis of cobalt-induced interfacial degradation. At the high temperatures present in typical CVD growth environments, diffusion of the binder phase into the film and subsequent diffusion of carbon from the film into the binder phase may occur. Given time, the diamond-substrate interface may be dissolved, thereby reducing the interfacial contact area and adhesion strength. The SEM image of a film that failed by this mode (a) was taken of the interface side of a diamond film following delamination. The pores displayed in the image were created in the diamond film by binder phase diffusion from the bulk into the film during deposition. Within the interior of the voids, small spheres of Co have solidified upon cooling down from deposition temperature. The corresponding energy dispersive spectrum (b) identifies the composition of the small, bright spheres as primarily cobalt. Co-diamond interactions may also be viewed in cross section as observed in the composite SEM image of a diamond WC-Co interface shown in (c).
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This SEM image was taken of the interface region between a diamond film and WC-Co substrate in which the interface between the diamond and substrate has resulted in the diffusion of carbon from the diamond into the binder phase of the substrate. Upon cooling down from deposition, the carbon in the binder phase has precipitated out in the form of graphite (shown as the ‘‘fuzzy’’ dark contrast material in the central region of the image). (From Ref. 74.)
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common grades of etched WC-Co in applications that generate lower cutting forces such as nonmetallic machining. b. Diffusion Barrier Approaches. Recognizing that a physical barrier or so-called diffusion barrier to diamond-binder interaction may improve adhesion by preventing interaction between the binder phase and the diamond film, other researchers [70–72] have developed techniques for the
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formation of interlayers between the CVD diamond film and WC-Co substrate. Proper selection of such a layer could also reduce residual stresses between the diamond film and the underlying substrate by choosing an interlayer material with a coefficient of thermal expansion that falls between the film and underlying substrate. Destructive analysis of commercial CVD diamond products produced by various industrial vendors indicates that interlayer technology is used; however, the specific intent of the interlayer design was not evident and observed failures were typically at the interlayer.* Testing indicated that the specific approaches studied in commercial samples failed to yield sufficient interfacial toughness to resist delamination in abrasive aluminum machining; however, the interfacial integrity of the same products may have been suitable for other less aggressive operations. Academic research on interlayers reported by Kupp [71] revealed a suppression of chemical interaction but generally poor cutting test performance based on industrial expectations. Although diffusion barrier approaches can add measurable cost and complexity to adhesion treatment, it is quite possible that unique interlayer solutions will evolve as the CVD diamond industry advances. c. Mechanical Roughening Techniques. Other researchers have recognized that rough interfaces potentially provide an increase in the adhesion strength of CVD diamond films or WC-Co materials. This has led to methods that include mechanical roughening by abrasion, wet-chemical etching of the WC phase [73], and laser ablation or blasting with abrasive grit. Unfortunately, most of these techniques form an interface that has sustained significant damage and actually reduce the interfacial toughness rather than improve it. d. High-Temperature Heat Treatments. Saijo and coworkers [73] discuss a process for treatment of WC-Co with a binder phase of 4 wt% or less using a decarburizing gas composed of oxygen and hydrogen at a temperature between 500 and 1200⬚C. Although the decarburization of the freesurface WC grains promotes chemical bonding between the diamond film and substrate through recarburization during CVD diamond deposition, the method produces a free surface in which the WC grains are smaller than in the bulk. According to Olson and Windischmann [74], this process therefore does not provide the crack deflection or interfacial toughening mechanism believed to be essential in highly abrasive applications and may create an embrittled surface layer due to decreased grain contiguity.
*Comment is based on studies performed within Norton Diamond Film prior to business closure.
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e. Present Status and Future Needs for Adhesion Treatments. The often undisclosed, proprietary adhesion treatments used by diamond coating vendors have been successful in producing suitable adhesion for selected applications but have still not resulted in diamond-coated WC-Co products that can be used for the full range of machining applications envisioned for CVD diamond a decade ago. For instance, diamond-coated tungsten carbide products offer measurable performance advantages in machining selected nonmetallic materials; however, adhesion issues have clearly limited the success in machining common metallic workpiece materials such as aluminum. It now seems apparent that the nature and extent of adhesion challenges for diamond-coated tungsten carbide cutting tools and wear parts create the need for a range of adhesion treatment solutions that are selected based on the specific demands of the application. For purposes of instruction, Sec. III.B.3 contains a case study involving adhesion treatment technology developed specifically for metalcutting insert manufacture. Section IV.B.1.a contains related metal cutting results. 3.
Case Study: Development of an Adhesion Treatment for Metalcutting Inserts
In order to maximize adhesion strength of diamond coatings of WC-Co substrate materials, it is essential to optimize many key features of the diamond-substrate interface. Unless all of these characteristics are in place, optimal adhesion strength cannot be realized for the most abrasive applications. In an approach reported by Olson and Windischmann [74], the essential characteristics of the diamond-substrate interface are produced that, acting together, optimize the adhesion strength of the diamond film to the underlying substrate. These characteristics fall in three general categories: (1) chemical composition, (2) phase composition, and (3) microstructural composition. The remainder of this section contains a detailed discussion of the specific attributes of this adhesion treatment for WC-Co cutting tool inserts. The chemical composition of the interface is altered to avoid deleterious binder phase–diamond reactions that can reduce the chemical bonding of the diamond film to the substrate and also induce a phase transformation of the diamond film to graphite. However, unlike techniques that remove the binder phase to some depth below the exposed WC-Co substrate surface, binder phase removal is done in a way which limits removal to an area that is directly exposed to the CVD growth species or ‘‘free surface.’’ In addition, the composition of the WC phase at the free surface is altered to maximize the density of direct chemical bonding between the diamond film and substrate. Unlike chemical etching methods which decarburize the WC grains
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by chemically ‘‘attacking’’ them, this is performed in a manner intended to maintain the mechanical properties of the substrate and interface. Finally, the microstructural composition of the interface is altered to minimize crack nucleation sites due to interfacial voids and provide a toughening, crack deflection mechanism that resists interfacial crack propagation. This feature essentially arrests or deflects cracks that may nucleate at the interface and impedes the propagation of these cracks, thereby suppressing delamination. The process utilizes commercially available WC-Co cutting tool inserts with a specific grain size and composition. The typical composition of the starting grade is 94.0 wt% WC, 6.0 wt% Co with an average grain size of the WC phase of about 1–2 m (Sandvik H13A, Carboloy 883, and others). Because the substrate in this case is in the form of a commercially available cutting tool insert, the free surface of the material is usually finish ground to final dimensional tolerances. Therefore, the free-surface usually displays the impression of residual grind lines and surface damage shown in Fig. 19a. This surface contains fragments of WC grains that have been fractured during the grinding step and a binder phase, which is composed mostly of Co, smeared across the surface. In the current approach, the free-surface binder phase is vaporized using a proprietary technique. Because of the nature of the process, the freesurface WC grain size is increased while the stoichiometry of the WC grains is shifted to a slightly C-deficient ratio. This leaves the free surface or ‘‘shell’’ of individual WC grains at the free surface of the substrate with a slightly decarburized ‘‘case.’’ Figure 19b shows the marked change in surface structure relative to the as-received WC-Co surface in Fig. 19a. Elemental analysis of the adhesion treated surface (Fig. 19b) by EDS indicates a surface that is essentially free of metallic binder phases (Fig. 19a). The combination of the simultaneous vaporization of the binder phase, which consequently increases the surface free energy of the WC grains, and the slight decarburization of the WC grains increases the dangling bond density of the free-surface grains. Enlargement of these free-surface WC grains is induced simultaneously with the binder phase vaporization and shift in stoichiometry. Thus, all three components of the required interface may be evolved: (1) the binder phase is removed through vaporization, (2) the chemical bonding between the diamond film and substrate may be maximized through a recarburization/diamond film growth step, and (3) the enlarged WC grains coupled with a high nucleation density provide the crack deflection enhancing microstructure that toughens the interface. All of these features are created using a one-step process developed to produce these characteristics in a way that minimizes the trade-offs of other important physical features of the carbide substrate.
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Figure 19 (a) SEM micrograph showing the finish-ground surface of a 6% CoWC cutting tool. Most adhesion technologies are targeted at removal of the metal binder that is smeared on the surface during grinding. (b) Using the surface modification technique described in Sec. III.B.3, the binder phase can be selectively removed from commercially available grades of as-ground WC-Co cutting tool inserts. As the image depicts, the WC grains are restructured to create a pristine WC surface ready for diamond coating. EDS spectra (not shown) indicate that an essentially cobalt-free substrate surface can be produced by this technique. (From Ref. 74.)
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Following the evolution of the surface microstructure and chemical composition of the substrate as just described, the surface is coated with a diamond film [74]. Diamond growth was carried out in both high-rate, microwave plasma CVD and DC arcjet CVD systems under high deposition rate conditions that suppress bulk-to-interface diffusion of the binder phase. The early stages of the CVD diamond process were tailored to optimize the nucleation density without formation of a nondiamond interphase. The characteristics of the diamond film have been developed to optimize the wear properties for abrasive environments as described in Sec. III.A and previously by Olson and Dawes [28]. The interface structure of a diamond-coated WC-Co cutting tool insert prepared by this technique is contrasted with that of a diamond-coated insert prepared by chemical etching in Fig. 20. Section IV.B.1.a contains the aluminum machining performance of inserts produced using this specific interfacial engineering approach.
IV.
CVD DIAMOND CUTTING TOOL PERFORMANCE
As discussed in Secs. II and III, there are many factors that can affect the performance of a specific CVD diamond cutting tool design in a chosen application. This section contains application case studies of both thin-film and thick-film CVD diamond cutting tools. The case studies are intended to clearly show that the relative performance advantage of CVD diamond tooling over conventional tooling varies widely depending on the specifics of the machining application and the tool design concepts employed. The case studies include discussion of the most critical application characteristics as well as the cutting tool wear mechanisms that dictate tool life and machining efficiency. A.
Freestanding Thick-Film CVD Diamond Cutting Tool Performance
1.
Thick-Film CVD Diamond Insert Performance
a. ‘‘High-Silicon’’ Aluminum Machining with Thick-Film CVD Diamond. As noted in Table 2 (Sec. II.C.2), silicon is commonly added to aluminum compositions to modify alloy properties. The Si-Al alloy has a eutectic near 12.2% Si-Al. Higher amounts of silicon in the composition result in the precipitation of hard, abrasive silicon particles in the workpiece that enhance the wear resistance of the aluminum surface. For instance, A390 (18% Si-Al) hypereutectic aluminum, which is commonly used in automotive pistons, contains approximately 6% hard silicon particles that vary in
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Figure 20 Interfacial comparison of two adhesion treatment approaches. Suitable diamond deposition onto substrates produced using the surface modification technique described in Sec. III.B.3 yields an interface that is free of large voids and of the chemical interfacial degradation observed in Co-diamond interactions. (a) Crosssectional SEM image of the diamond film–substrate interface. (b) For comparison, a cross-sectional SEM image of a diamond film on an etched WC-Co substrate. Note the porosity of the binder-depleted zone (embrittled layer) and the low interfacial roughness of the interface shown in (b) versus (a). (From Ref. 74.)
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size and distribution depending on the aluminum processing technique. According to Whitacre, a piston industry representative [75], ‘‘diamond tooling is an absolute requirement to successfully machine hypereutectic [high-silicon] alloys.’’ Because the volume of silicon particles increases with increasing silicon content, the machinability of low-silicon hypoeutectic aluminum alloys is significantly better than that of high-silicon hypereutectic aluminum alloys. Although PCD tooling is widely used for hypereutectic alloy machining in the piston industry, some companies have been benefiting from the enhanced tool life offered by brazed-on, thick-film CVD diamond tips since the early 1990s. CVD diamond properties captured in a properly brazed and finish-ground thick-film CVD diamond tool tip can yield extended tool life in hypereutectic aluminum alloys while producing better and more consistent surface finish than PCD. For a piston machining operation, Hay [8] reported that a CVD diamond tool offered three times the tool life of PCD in a finish profile turning operation. Using a cutting speed of 700 m/min, water–soluble oil coolant, and a depth of cut of 0.15 mm, A390 hypereutectic aluminum pistons were machined in a two-pass operation with the first pass using a feed rate of 0.30 mm/revolution and the second pass using a feed rate of 0.15 mm/rev. The 25-m grain size PCD tools yielded approximately 10,000 pistons/tool before finish specification was lost. By comparison, the properties of CVD diamond enabled the facility to produce an average of 30,000 pistons/tool in the same operation. The 100⫻ SEM image in Fig. 7 is that of a wear scar found on a CVD diamond piston turning tool after machining A390 aluminum pistons. This tool has worn to a point that requires regrinding of the flank face of the tool in order to resharpen the cutting edge. Depending on the extent of edge damage before each resharpening, tools in this application are reground more than five times before disposal. During its useful lifetime, a single CVD diamond tip is capable of machining over 100,000 high-silicon aluminum pistons. Wear of the CVD diamond can be seen along the cutting edge where the finish-ground flank face of the tool meets the polished rake face. The higher magnification 500⫻ image in Fig. 21 shows that the wear scar has some striations that may be the result of anisotropic wear of the diamond crystallites or, more likely, silicon particle impact damage. Hard particle damage to the polished rake face would tend to evolve into striations along the flank face wear surface, which is in constant contact with the rotating piston. Although the tool tip is visually reflective with no significant buildup of aluminum, the dark region of contrast at the top of the 100⫻ image in Fig. 7 is likely to be a submicrometer aluminum layer resulting from chip flow. Also noteworthy is the apparent waviness of the rake face in the bright region on the left side of the 100⫻ image. This feature is
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Figure 21 Flank-face SEM image of a freestanding CVD diamond cutting tool after use in turning high-silicon aluminum pistons. This image is a 500⫻ close-up of the 100⫻ image shown in Fig. 7. A single CVD diamond tip is capable of machining over 100,000 high-silicon aluminum pistons. (From Norton Diamond Film/Saint Gobain Industrial Ceramics, Inc.)
probably induced by the erosive nature of the Si-Al chips that are formed in this region of the tool. b. ‘‘Low-Silicon’’ Aluminum Machining with Thick-Film CVD Diamond. Although hypoeutectic low-silicon aluminum compositions have good machinability and can generally be machined with nondiamond tooling materials such as tungsten carbide, diamond-tipped tooling is commonly used in mass production environments, which benefit from diamond’s ability to endure many hours of stable, high-speed operation. Drury has reported (T. Drury, personal communication, 2000) [1] performance in an aluminumfinishing operation in which CVD diamond-tipped tooling not only outlasted PCD by 30% but also maintained the required surface finish throughout the tool life. Hay [8] has reported the results of a finish-boring application where hard, anodized 6061 aluminum turbine cases were machined with both CVD and PCD diamond-tipped tools. Using a cutting speed of over 457 m/min, a depth of cut of 0.013 mm, and a feed rate of 0.05 mm/rev, CVD diamond outperformed fine-grain PCD (5 m grain size) under stringent surface finish specifications. The CVD diamond tool reportedly increased tool life from 106 parts per tool up to 345 parts per tool and surface finish was more consistent than that of the PCD tooling. This is one of many examples where
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thick-film CVD diamond not only extends wear life but also produces and maintains superior surface finish relative to PCD tools. c. Metal Matrix Composite Machining with Thick-Film CVD Diamond. As noted in Table 2 (Sec. II.C.2), metal matrix composites (MMCs) are typically reinforced with over 20% ceramic particles or fibers, making them extremely difficult or impossible to machine with conventional tooling such as cemented carbide. One of the present uses of MMCs is for automotive brake rotors where aluminum matrix composites are employed. MMCs are also of interest for improved high-speed rail braking systems and components for the electronics industry. CVD diamond tooling has shown promise as an alternative to PCD tooling or, in some cases, grinding or water-jet cutting MMCs. In tests performed by an Austrian research group [76] using a 30% SiC reinforced aluminum-matrix composite, it was determined that larger grain size PCD significantly outperformed smaller grain size PCD. In the same tests, CVD diamond-tipped tools outperformed all forms of PCD tested in finishing the MMC material. As an example, after 10,000 m of finishing, the CVD diamond tool had a wear scar less than 25 m in width, whereas 25- and 40-m grain size PCD had worn to over 120 and 230 m, respectively. The superior wear resistance offered by CVD diamond is attributed to its binder-free, pure diamond construction, which reduces the tendency for diamond grain ‘‘pullout’’ when eroded by the SiC particles. Because CVD diamond can be manufactured with high diamond-to-diamond grain boundary strength, the grain size of the diamond crystallites is not nearly as critical for CVD as it is for cobalt-sintered PCD. However, as discussed in Sec. III.A, the grain boundary strength of CVD diamond deposits is likely to be very critical in abrasive applications such as MMC machining. The high erosion resistance of thick-film CVD diamond in machining MMCs and high-silicon aluminum materials is probably due to the combination of diamond’s hardness, fracture toughness, lubricity, and other key properties. High-silicon aluminum or MMC machining requires very high erosion resistance to hard particles. In general, the metal matrix of both MMCs (including high-silicon aluminum) is free-machining aluminum; therefore, tool wear is dictated by the size and distribution of the hard particulate reinforcement. Table 3 contains erosion resistance data for CVD diamond and other materials known for their hardness or wear resistance. In these tests, CVD diamond clearly showed superior erosion resistance compared with alternative tooling materials with a relative volume loss four times lower than for PCD and 120 times lower than for tungsten carbide. The performance of thin-film tooling in erosive workpiece materials requires not only erosion-resistant diamond but also interfacial impact resistance be-
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Table 3
Relative Erosion Resistance of CVD Diamond and Other Materials
Material
Relative volume loss
CVD diamond (low pressure, vapor deposited) PCD (high pressure, cobalt sintered) Cemented tungsten carbide (6% Co-WC sintered) Aluminum oxide ceramic (99.5% Al2O3) Silicon carbide ceramic Silicon nitride ceramic Erosion test parameters: 2% SiC particles in water (30 grit; >0.5 mm average particle size) 91.4 m/sec abrasive flow rate 45⬚ angle of impingement
1 4 120 220 360 920
Source: Norton Diamond Film, SGIC, Inc. internal test data.
tween the tooling material and the coating. The performance of diamondcoated inserts in high-silicon aluminum machining is discussed in Sec. IV.B.1.a. The performance of diamond-coated drills in MMCs is discussed in Sec. IV.B.2.c. d. Sintered Tungsten Carbide Boring with Thick-Film CVD Diamond. Tungsten carbide not only offers unique solutions as a cutting tool material as discussed in Secs. I.B.2 and II.B, but also is commonly used in tough, wear-resistant wear parts. In selected applications, tungsten carbide is precision machined using PCD, PCBN, and, more recently, CVD diamond. In a boring application, a 25% Co-WC cylinder with a 20 mm inner diameter and 40 mm length was machined using a speed of 30 m/min, feed of 0.05 mm/rotation, and depth of cut of 0.12 mm. Flood coolant was used. The results are summarized in Fig. 22. The CVD diamond thick-film tool machined eight parts before failure. PCD machined only five parts per tool and PCBN was only able to machine one part. Another carbide boring application demonstrated the ability of CVD diamond to produce a similar finish to single-crystal diamond tooling. The ability to machine dry, with no coolant, has also been demonstrated with CVD diamond thick-film tooling. 2.
Thick-Film CVD Diamond Rotary Tool Performance
The performance advantages of thick-film CVD diamond relative to PCD are similar to that observed for thick-film CVD diamond inserts. For example, in a 19% Si-Al wrist pin bore reaming operation, a PCD reamer was replaced with a thick-film CVD diamond blade reamer insert. Reamer life
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Figure 22 Performance comparison of CVD diamond, PCD, and PCBN in sintered tungsten carbide boring (25% Co-WC). (Data provided by a customer of Norton Diamond Film/Saint Gobain Industrial Ceramics, Inc. and Ceratonia.)
was extended by three to four times using a cutting speed of 628 m/min, speed of 51 mm/rev, with a depth of cut of 0.076 mm per side. Even though the potential performance advantage are significant, the present use of thickfilm CVD diamond in rotary tooling including reamers, end mills, routers, and drills is much less extensive than the usage in the simpler geometries of inserts. One of the major barriers for thick-film CVD rotary tools is the inability to EDM ‘‘wire-cut’’ CVD in a manner similar to that with PCD. Many of these challenges will ultimately be overcome as methods of fabricating with flat, thick-film CVD shapes evolve or alternative, adherent solutions are developed with thin-film diamond. Since diamond-coated tungsten carbide rotary tools are the present focus of the industry, performance advantages of thin-film diamond rotary tool designs are discussed extensively in Sec. IV.B.2. B.
Thin-Film CVD Diamond Cutting Tool Performance
1.
Thin-Film CVD Diamond Insert Performance
Depending on the application, diamond-coated inserts can show a cost-performance advantage relative to uncoated carbide, other hard coatings, and
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PCD. Although there are cases where single corners of a CVD diamond insert will outperform a single PCD tip, the indexability (and disposable design) of diamond-coated inserts is generally a key advantage relative to regrindable PCD tips. CVD diamond-coated inserts generally cannot produce the superior surface finish that is expected from fabricated PCD tips. In some cases, diamond-coated insert design limits the use of diamond-coated tools to roughing and intermediate finishing processes. In such instances, fabricated, thick-film CVD diamond tips are a viable alternative for long wear life and improved surface finish over PCD. a. Case Study: Performance of Diamond-Coated Inserts. To evaluate fully the level of performance of the CVD diamond-coated WC-Co cutting tool inserts manufactured using the process described in Sec. III.B.3, inserts were tested in a range of industrial applications under a variety of conditions involving workpiece materials ranging from graphite to hypereutectic aluminum. The tests were carried out in machining laboratory settings as well as various industrial settings by independent evaluators, and the results were compared with those for tools traditionally used in the specific application. HIGH-SILICON ALUMINUM TURNING. In A390 high-silicon aluminum turning tests, the tools were demonstrated to last significantly longer than uncoated carbide tools or PCD-tipped tools. Using a cutting speed of 680 m/min, feed rate of 0.2 mm/rev, and depth of cut of 1.0 mm, diamondcoated C2 grade WC-Co TPG 321 style inserts were tested against uncoated carbide, coarse grain PCD (25 m grain size), and thick-film CVD diamond of the same tool geometry. Both the PCD and thick-film CVD tools are regrindable by design; the diamond-coated and uncoated carbide tools are disposable. Tools were tested to a predetermined failure criterion defined by 0.375 mm of flank wear. Figure 23a shows the results of the tool life testing. Thick-film CVD diamond (500 m diamond brazed-on and finish ground) produced the best per tip tool life, followed by thin-film CVD diamond (25– 35 m coating), PCD, and uncoated C2 grade, 6% Co-WC carbide. Because diamond-coated tools are indexable and a TPG 321 tool is triangular, the per tool lifetime was actually highest for the disposable CVD diamond thinfilm tools. The thin-film tool’s rough interface (created by the adhesion treatment discussed in Sec. III.B.3) coupled with the faceted diamond morphology critical to high wear resistance in the A390 workpiece limits the ability of this tool design in finishing selected workpiece materials. As Fig. 23b summarizes, the finish-ground cutting edges on both PCD and thickfilm CVD diamond yielded surface roughness values near 0.4 m R a, whereas, the diamond-coated tools produced values near 0.75 m R a. Using this and other case studies, the adhesion treatment approach discussed in Sec. III.B.3 proved to offer suitable adhesion to endure various
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Figure 23 Comparison of PCD, CVD diamond thick-film, CVD diamond thinfilm, and uncoated WC-Co tooling in turning high-silicon aluminum. Tool lifetime is compared in (a). The surface finish of the Mahle A390 (⬃18% Si-Al) workpiece is shown in (b). (From Norton Diamond Film/Saint Gobain Industrial Ceramics, Inc.)
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abrasive wear and impact-oriented applications. Unfortunately, the efforts to engineer the necessary properties into the diamond-coated WC-Co ‘‘composite’’ required a trade-off in surface finish of the workpiece that generally limits the use of this particular tooling design concept to intermediate finishing and roughing. Figure 6 (Sec. II.D.1.a) depicts this surface roughness versus adhesion trade-off. Figure 6b shows a relatively rough, diamondcoated tool produced by this method. A uniform wear pattern can be observed at the cutting corner. Figure 6a shows a relatively smooth, but delaminated tool produced by a wet chemical etching adhesion treatment. As Fig. 23 depicts, the thick-film, brazed-on diamond tip not only demonstrated superior life relative to coarse-grained PCD but also delivered comparable finish. b. Performance of Other Thin-Film Diamond Tool Designs HIGH-SILICON ALUMINUM TURNING. A CVD diamond machining study by Oles et al. [23] reported turning test results for two hypereutectic alloys (Reynolds A390 and Mahle 138) that contained nominally 18% silicon. Machining conditions were as follows: 762 m/min, 0.127 mm/rev, 0.635 mm depth of cut, 15⬚ lead angle, and flood coolant. The tested tools were SPGN120308-style inserts with a diamond-coating thickness of approximately 30 m. Regardless of the alloy type, diamond-coated and PCD tools lasted approximately two orders of magnitude longer than the uncoated carbide tools, which reached the failure criterion of the testing within seconds. This exceptionally rapid wear was induced by the hard silicon particles and clarifies why diamond tooling is generally accepted as the only suitable solution for high-productivity machining of high-silicon aluminum alloys [75]. The study also showed the strong influence of the workpiece microstructure on the wear rate and general performance of CVD diamond–coated tools relative to PCD. In Mahle 138 alloy, the steady-state wear rates on the coating average 0.005 mm/min, whereas the wear rate averaged 0.012 mm/min when machining Reynolds A390. The higher wear rates observed in the Reynolds A390 were attributed to the higher level of impact damage induced by the larger, less homogeneous silicon particulates. The microstructure of the A390 alloy contained a nonuniform distribution of relatively larger, irregularshaped silicon particles (⬃0.09 mm size, 100 particles/cm2), whereas the 138 alloy contained a uniform distribution of smaller, spheroidal particles (⬃0.02 mm size, 160 particles/cm2). LOW-SILICON ALUMINUM TURNING. Oles et al. [23] also reported turning tests using CVD diamond-coated tools in hypoeutectic 383.2 aluminum (11% silicon) using the same machining conditions summarized before for the hypereutectic alloy study. The steady-state diamond wear rate
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was more than five times lower than the wear rates reported for the highsilicon alloys in the previous section. The wear rate for uncoated carbide was 0.011 mm/min compared with 0.001 mm/min estimated for the diamond-coated tooling. The authors noted that the much lower wear rates of the cemented carbide tooling explain the general acceptance of uncoated and TiN-coated carbides as a solution for machining hypoeutectic, low-silicon aluminum alloys. CARBON-FILLED PHENOLIC COMPOSITE TURNING. Hay [8] has reported the performance of diamond-coated silicon nitride ceramic tools in a 30 vol% carbon-filled phenolic resin machining application. Precision parts were turned using a cutting speed of 122 m/min, feed rate of 0.003 ipr, and a depth of cut of 0.635 mm. To counter the corrosive effects of the workpiece, aluminum oxide ceramic tools were used to attain approximately 50 parts per cutting corner and a marked cost-lifetime advantage over metalbonded PCD and tungsten carbide tooling. The introduction of diamondcoated silicon nitride tools into the same operation now offers up to 650 parts per corner due to an unmatched combination of corrosive and abrasive wear resistance. 2.
Thin-Film CVD Diamond Rotary Tool Performance
a. General Status of Nonmetal Machining with Diamond-Coated Rotary Tools. Some of the most impressive tool life improvements with CVD diamond have been reported in applications using diamond-coated rotary tools. In cases where the level of adhesion is suitable for the application, the lifetime advantage of a diamond-coated tungsten carbide tool is often dramatic. Because of the geometric complexity of most rotary tools, PCD (or thick-film diamond) tooling is generally not a cost-effective option; therefore, tungsten carbide and hard coatings on carbide are commonly used in many high-throughput environments. Lifetime advantages of diamondcoated tools relative to other hard coating or uncoated carbide generally fall in the range of 10 to 100 times for machining nonmetallic materials including structural composites, abrasive or corrosive polymers, green ceramics, graphite, and hard carbon. Selected case studies of nonmetallic machining applications are summarized in Sec. IV.B.2.c. b. General Status of Metal Machining with Diamond-Coated Rotary Tools. Although similarly impressive performance advantages have been observed in metal machining applications, not all vendors recommend their diamond-coated rotary tool products for metalcutting at the time of publication. Although much of the potential for diamond-coated tooling is projected to be automotive and aluminum machining, the present use of diamond-coated rotary tools in metalcutting applications is not nearly as
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widespread as for nonmetals. Engineering diamond-coated tools to sustain the shear forces of metal machining has been a major barrier to expansion of their application range. There have been cases where commercial claims of metalcutting capability have been retracted from the marketplace because of reproducibility issues in the field. Even with these challenges, rotary tool users who machine nonferrous metals such as aluminum alloys, copper, brass, bronze, and magnesium should closely monitor the expanding application range of CVD diamond-coated tooling as commercial adhesion technology and application know-how continue to evolve. The following section contains a summary of a case study on MMC drilling. c. Diamond-Coated Rotary Tool Case Studies GRAPHITE AND HARD-CARBON MATCHING. The advantages of CVD diamond-coated rotary tools in enhancing graphite-machining operations are presently aggressively marketed by many vendors and are becoming well known in the machining industry [18–20,77]. Ten to 20 times tool lifetime advantages are commonly observed in small job shops and large machine shops involved in EDM sinker-die electrode and other graphite component production. Success has been achieved with both outdated milling machines and high-speed machining centers. High-speed machines amplify the advantages of CVD diamond as other hard coatings and uncoated carbide both fall short at operating speeds in excess of 900 m/min. When forms of carbon harder than graphite are machined, the relative performance numbers are often even more dramatic than the graphite machining results. In an antimony-impregnated hard-carbon (carbon-graphite) end milling application (T. Uemura, personal communication of field test data, 1999), a diamond-coated end mill with a faceted 10 to 15-m coating was used as a replacement for a commercially available diamond-like carbon (DLC) coating. Using a 9.5 mm diameter, four-flute, square end mill, notches were machined into a gas seal ring. DLC-coated end mills were able to produce approximately 120 parts per tool using the following conditions: 1000 rpm (30 m/min), 38 mm/min feed rate, and a depth of cut of 5.3 mm. Not only did CVD diamond demonstrate superior performance at the same conditions, but also the properties of the coating enabled operation under more aggressive conditions: 3000 rpm (a threefold increase), 1016 mm/min feed rate (a 26.7-fold increase), and a depth of cut of 5.3 mm. Under these conditions, the CVD diamond-coated end mill produced approximately 11,000 parts and the customer realized tool life increases of over 90 times relative to the DLC-coated tool. More important, the machining process productivity was dramatically affected by the increases in speed and feed enabled by the diamond-coated tool as per part cycle times were decreased from 20 to 3 minutes per seal ring.
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Although the present costs of DLC tooling are markedly lower than those of CVD diamond coatings, the cost-performance advantage of pure, crystalline, CVD diamond is clear in many applications such as the preceding hard-carbon case study. Most DLC coatings are limited in higher temperature cutting applications (often induced by higher speed operations) because of phase instability at temperatures as low as 250–350⬚C. The level of phase instability has been determined to be a function of the hydrogen content in DLC materials with pure amorphous carbon having the best stability of materials in the continuum of amorphous hydrogenated carbon (aH:C) materials. Although Horsfall [78] has reported properties of noncrystalline carbon coatings that have low hydrogen content and high-temperature stability above 500⬚C, the continuum of diamond-like carbon materials are limited in performance relative to CVD diamond because of their relative thermal instability. Furthermore, due to compressive stresses in DLC coatings, they are often limited to a coating thickness of approximately 2 m, over five times thinner than many commercially available diamond-coated tools. In addition to per tool cost advantages, DLC coatings have an advantage over diamond coatings in edge sharpness (a function of coating thickness), coating surface roughness (typical of amorphous coatings), and much lower sensitivity to substrate selection (due to low-temperature coating process). PRESINTERED CERAMIC MACHINING. In the presintered ceramic machining test discussed earlier in Sec. II.C.1, the performance of 9.5 mm diameter, four-flute diamond-coated end mills was compared with that of uncoated tungsten carbide and PVD TiN-coated carbide tools of the same geometry. The spindle speed was 10,000 rpm, with a feed rate of 508 mm/ min (20 ipm) corresponding to a cutting speed near 300 m/min (1000 sfm) and a chip load of approximately 0.012 mm. The presintered ceramic workpiece was machined with no coolant using alternating, 7.6 cm long climb (up-cut) and conventional (down-cut) passes with a 4.8 mm radial depth of cut and a 12.7 mm axial depth of cut. Cutting forces on the workpiece were measured using a three-axis dynamometer and a computerized data collection system. Testing details have been reported elsewhere by Tanikella et al. [32] and Cline [31]. Figure 5 visually depicted the difference in cutting temperatures between diamond and the alternative tools in this test. This dramatic result has been attributed to the superior lubricity (low coefficient of friction) of the diamond coating when in contact with the ceramic workpiece material. According to Quinto et al. [79,80], the hardness of PVD TiN-coated carbide at 1000⬚C is approximately 25% of its room temperature hardness. This ‘‘hot hardness’’ limitation explains the lack of any performance advantage afforded by the PVD TiN coating relative to uncoated carbide under these
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cutting conditions. The reduced friction and wear were clearly reflected in the force and power data summarized in Table 4. In the first 7.6 cm long machining pass, the spindle power consumption for the TiN-coated carbide and uncoated carbide tools was over 10 times higher than the power consumption for the diamond tool after 80 passes. The lower cutting forces seen with CVD diamond-coated tools offer the potential for reduced yield losses in green or presintered ceramic machining. The low cutting temperature seen with the CVD diamond tool is viewed not only as an indication of the tremendous potential for CVD diamond tooling in high-speed machining of ceramics but also as the promise of CVD diamond coatings for machining many other difficult-to-machine abrasive materials. CARBON-EPOXY COMPOSITE MACHINING. Carbon-epoxy composites are of interest for lightweight, structural applications ranging from the aerospace industry to sporting goods. Because the hard, stiff carbon fibers in carbon-epoxy induce excessive wear in traditional cutting tool materials, some manufacturers have resorted to using alternative shaping techniques such as water jet cutting or computer numerical control (CNC) grinding with diamond grit tools. Other manufacturers are beginning to benefit from the advantages of CVD diamond-coated rotary tools in this difficult-to-machine material. Tanikella et al. [32] and Cline [2] have reported the results of carbonepoxy machining tests to determine the advantage of diamond-coated rotary tools over uncoated tungsten carbide two-flute, 9.5 mm diameter routers.
Table 4 Cutting Forces and Spindle Power Measured During Dry Machining of Presintered Ceramic Pass number
Fx (newtons)
Fy (newtons)
Fz (newtons)
Power (watts)
Typical measurements for uncoated or TiN-coated tungsten carbide end mills First pass (climb) Second pass (conventional)
200–250 300–400
100–150 500–700
20–50 80–100
800 1250–1400
Typical measurements for diamond-coated tungsten carbide end mills Average for pass 1 through 8 Average for pass 81 through 88 Source: Ref. 31.
40
40
15–25
50
60
60
25
50
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Using a spindle speed of 20,000 rpm with a table speed of 2540 mm/min, alternating climb milling and conventional milling passes were performed on a 17.8 cm long, rigidly fixtured panel. The selected cutting conditions were in accordance with the interests of a large commercial aircraft manufacturer. As shown in Fig. 24a, severe wear was observed on the carbide tool in less than 36 cm of cutting. Figure 24b shows that cutting forces for the uncoated carbide tool increased dramatically during the first (climb) machining pass. After 20 passes, the CVD diamond-coated tool was still sharp and running predictably as depicted in Fig. 24c and d. Cutting noise and cutting odor typical of uncoated carbide tools were dramatically reduced with the use of CVD diamond-coated routers. Figure 25 shows the machined cross section of the carbon-epoxy composite after 40 passes using a diamond-coated router. The machined surface is burr free and cut cleanly enough to enable observation of the layered structure of the composite. It is interesting to note that the spacings between the composite layers in Fig. 25 match the spacings between the worn notches in the uncoated carbide tool shown in Fig. 24a. CVD diamond-coated tooling clearly has the potential to minimize costly deburring operations in composites fabrication. Even with flood coolant, the hardness and frictional heating generated by the carbon fibers clearly exceeded the performance limits of the uncoated carbide cutting edge. In another carbon-epoxy machining study [81], production testing of drills, drill countersinks, and two-flute helical end mills in carbon fiber composites showed that CVD diamond-coated cutting tools last up to 50 times longer than tungsten carbide and more than twice as long as PCD tools. This same reference reported that CVD diamond-coated tools were projected to increase the lifetime of machine tools by as much as 50%, significant cost savings to a machine shop. FIBER-REINFORCED POLYMER MACHINING. Matching tests were performed in a glass fiber-reinforced polymer (GFRP) using CVD diamondcoated tungsten carbide and uncoated carbide. The specific GFRP studied contained nominally 65% glass fibers and is, therefore, extremely difficult to machine. SEM micrographs of the composite cross section shown in Fig. 26 depict the extremely high density of fibers having a nominal diameter of 10 m. Using 6.4 mm diameter, four-flute, square-end end mills, the GFRP was machined with no coolant at 12,000 rpm (239 m/min), a feed rate of 2438 mm/min, and an infeed of 3.2 mm. The selected machining conditions create a feed per tooth or ‘‘chip load’’ of 0.051 mm per cutting edge. Table 5 contains end mill diameter measurements for the CVD diamond-coated tool relative to an uncoated carbide tool. The dimensional stability of the CVD diamond-coated tool is dramatically better than that of the uncoated carbide tool. According to laser diameter measurements, the
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Figure 24 Comparison of uncoated and CVD diamond-coated tungsten carbide router performance in carbon-epoxy machining. (a) After only two cutting passes less than 36 cm, optical inspection of the uncoated carbide tool showed severe tool wear. (b) The rapid degradation of the uncoated cutting edge was further exemplified in the spindle power data for the first pass (made in the climb direction). (c) In contrast, the CVD diamond-coated carbide tool exhibited relatively insignificant amounts of wear after 40 cutting passes (711 cm.). (d) The operational stability of the CVD diamond-coated tool is shown in the spindle power data for the 17th through 20th passes. Note: 1 division = 0.5 mm in (a) and (c). (From Ref. 2.)
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Figure 24
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Continued
diameter of the uncoated carbide was reduced by over 40 m after only 62.9 cm3 of material removed, whereas the diamond-coated end mill lost less than 2 m of diameter after removing 20 times more workpiece volume (1265.5 cm3) under the same cutting conditions. As Fig. 27 shows, structural analysis of the wear scar verifies that the CVD diamond coating shows only
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Figure 25 Oblique view of the cross section of a carbon-epoxy panel that was machined in the 40th cutting pass with a CVD diamond-coated helical router. The layer spacings of the composite material seen in this image match the spacing between the worn notches in the uncoated carbide tool shown in Fig. 24a. (From Norton Diamond Film/Saint Gobain Industrial Ceramics, Inc.)
superficial wear of the faceted diamond peaks at this point of testing. In contrast, the uncoated tool has significant amounts of wear after machining less than 5% of the workpiece volume (62.9 cm3). The workpiece finish produced by the CVD diamond-coated tool was superior and did not deteriorate during the testing. The test was stopped when the uncoated carbide tool reached a point at which it had clearly worn out (62.9 cm3 of material removed). Table 6 contains data for the power consumption by the machine tool spindle observed during the testing. The data confirm that the dimensional stability of the CVD diamond-coated cutting tool results in a consistent power draw on the spindle, whereas, the power draw on the spindle had increased by 50–70% for the uncoated carbide tool. This excess energy is primarily dissipated as frictional energy during cutting, which tends to overheat both the cutting tool and the polymer-based workpiece. The machinist noted that the uncoated carbide tools were ‘‘too hot to touch’’ from the onset of the test and that they started to glow ‘‘red hot’’ after about 47 cm3 of machining. The CVD diamond-coated tools ran ‘‘cool to the touch’’ throughout the entire test, presumably because of the high lubricity of the diamond when in contact with the high density of glass fibers in the workpiece.
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Figure 26 SEM images of the cross section of a glass fiber-reinforced polymer (GFRP) component that was machined with a CVD diamond-coated end mill. The image in (a) depicts the high fiber volume percentage characteristic of the most abrasive GFRP materials. The image in (b) shows that the ⬃10 m diameter fibers are pulverized during the machining action of the diamond tool and clarifies the need for abrasion-resistant cutting tool material in machining this particular workpiece. (From Norton Diamond Film/Saint Gobain Industrial Ceramics, Inc. Test results and photographs courtesy of Walter Swann.)
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Table 5
Laser Measurements of End Mill Diameter from GFRP Composite Machining Study
Tool type Uncoated WC-Co CVD diamond–coated WC-Co
Volume of material removed (cm3) 62.9 1262.5 (20 times more)
Pretest tool diameter measurements (in.)
Posttest tool diameter measurements (in.)
Change in tool diameter due to tool wear
1 2 1 2
1 2 1 2
1 2 1 2
= = = =
0.24859 0.24862 0.24963 0.24970
= = = =
0.24692 0.24694 0.24960 0.24965
= = = =
⫺0.00167 ⫺0.00168 ⫺0.00003 ⫺0.00005
in. in. in. in.
(⫺42 m) (⫺43 m) (⫺0.76 m) (⫺1.27 m)
Source: Norton Diamond Film/Saint Gobain Industrial Ceramics, Inc. Internal test data provided courtesy of Walter T. Swann.
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Figure 27 Performance of CVD diamond-coated carbide routers in GFRP machining. The lifetime advantage of CVD diamond-coated WC-Co end mills relative to uncoated WC-Co was estimated to exceed 20 times based on laboratory testing. Low-magnification optical images of a diamond-coated WC-Co end mill (a) before testing and (b) after 1262.5 cm3 of GFRP material was machined indicated that no measurable wear was present. Low-magnification optical images of the uncoated WC-Co end mill (c) before testing and (d) after only 62.9 cm3 of GFRP material was machined indicated that extensive wear was present on the cutting edge. SEM images at 100⫻ magnification of the cutting edge taken after testing verified that essentially no wear was observed on the CVD diamond-coated end mill after 1262.5 cm3 of machining (e), whereas the uncoated tungsten carbide tool had suffered extensive abrasive wear after machining 20 times less (62.9 cm3) workpiece material (f ). (From Norton Diamond Film/Saint Gobain Industrial Ceramics, Inc. Test results and photographs courtesy of Walter Swann.)
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Figure 27
Continued
Consequently, a pungent burning odor of the polymer was noted only for the uncoated tools. The photographs in Fig. 26 are actually the machined surface created with a diamond-coated tool. Clearly, these conditions require the cutting edge to pulverize the hard, abrasive glass fibers during the machining process. The coating hardness and interfacial impact resistance are both very critical in ensuring predictable, abrasive wear with the CVD diamond-coated tool. POLYCARBONATE MACHINING. In a polycarbonate milling application, a small job shop replaced a tungsten carbide rotary tool with a 12.7 mm diameter CVD diamond-coated end mill (coating thickness of 10–15 m). Using an outdated milling machine, the user was able to machine over 3000 polycarbonate parts using the CVD diamond-coated end mill. By com-
Table 6
Spindle Power Consumption During GFRP Composite End Milling Spindle power consumption (watts) End of testing
Start of testing Machining direction → Uncoated WC-Co CVD diamond-coated WC-Co
Climb cutting
Conventional cutting
Climb cutting
Conventional cutting
80 75
80 80
120 80
135 80
Source: Norton Diamond Film/Saint Gobain Industrial Ceramics, Inc. Internal test data provided courtesy of Walter T. Swann.
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parison, an uncoated tungsten carbide end mill typically produced 30–35 parts per tool. The machine-limited cutting conditions were 3200 rpm with a feed rate up to 280 mm/min. The 100 times lifetime advantage of the CVD diamond tool relative to WC-Co is attributed to the corrosion resistance of the pure diamond composition relative to the cobalt-sintered tungsten carbide (discussed in Sec. II.A.8). If the binder phase of WC-Co (or PCD) is eroded, the wear mechanism is likely to shift from a slow, abrasive wear of the hard particle phase (WC or diamond, respectively) to a rapid, corrosioninduced disintegration (grain pullout) of the particulate composite at the cutting edge. MMC DRILLING. In magnesium-matrix composite drilling tests [82], CVD diamond coatings were shown to offer superior performance relative to uncoated tungsten carbide and TiAlN-coated tungsten carbide. TiAlN and CVD diamond coatings were being tested in an effort to find a lower total cost solution than PCD. The workpiece material compositions included particle-reinforced Mg (12 vol% SiC), fiber-reinforced Mg (20 vol% Al2O3), and hybrid-reinforced Mg (15 vol% SiC particles ⫹ 5 vol% Al2O3 fibers). The performance advantage of the CVD diamond-coated drills was most dramatic in the SiC-containing composites, where excessive wear was observed in TiAlN-coated and uncoated carbide. Although a dramatically improved performance was observed for the TiAlN coating in the Al2O3 fiberreinforced MMC, CVD diamond performed well in all composites tested. The satisfactory performance of TiAlN in the Al2O3 fiber-reinforced MMC, but not the SiC-containing composites, demonstrates the criticality of selecting the appropriate cutting tool material for each specific application. The superior performance of CVD diamond was attributed to several factors including diamond’s higher hardness relative to both SiC and Al2O3 as well as the higher level of adhesion of the diamond coatings relative to TiAlN coatings. Diamond’s unmatched hardness reduces the sensitivity to changes in the hardness of the MMC ceramic reinforcement phase. EVOLVING PROMISE FOR FERROUS MACHINING. Although many researchers believe that diamond tools are unable to machine ferrous materials because of the chemical incompatibility discussed in Sec. II.A.8, Kanda et al. [17] reported progress in the use of CVD diamond-coated rotary tools in machining ferrous materials. In tests using annealed gray cast iron (FC250; JIS G 5501), a diamond-coated, two-flute, ball-nose end mill (4 mm diameter, 8 mm flute length, and 80 mm total length) was tested. Tools were operated in both an up-cut (climb) and down-cut (conventional) machining direction using compressed air as a coolant under the following cutting conditions: 24,000 rpm, 2 m/min feed rate, with an axial depth of cut of 0.5 mm. Under these conditions, the wear of the CVD diamond-coated tool was less than 10% of that of an uncoated tool. In another test [17], the same
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researchers reported the relationship between the wear rate of the CVD diamond coating and the cutting conditions in carbon steel (S50C; JIS G 4051) with a hardness of 200 HB. Using a ball-nose end mill (2 mm diameter, 4 mm flute length, and 60 mm total length), the carbon steel was machined for a total cutting length of 4.0 m using various operating conditions: 8000 rpm and 16,000 rpm with an axial depth of cut of 1.5 mm and a radial infeed of 0.1 mm. Tests were performed using both compressed air and water-soluble emulsion. The work of Kanda et al. [17] indicates that the selection of a proper tool design and operating conditions can suppress the chemical corrosion of the diamond cutting edge, thus enabling superior performance over traditional tooling materials.
V.
OTHER MECHANICAL APPLICATIONS OF CVD DIAMOND
Both the unique characteristics of CVD diamond and the flexibility of CVD diamond processes allow the use of CVD diamond in mechanical applications beyond cutting tools. For decades, natural and synthetic (HPHT) single-crystal diamond has been used not only for specialty cutting applications as noted in Sec. I.C.1 but also for unique wear applications such as wire dies, engraving tools, and dressing tools for grinding wheels. The advent of thin-film and thick-film CVD diamond offers an alternative type of diamond for these applications, and the ability to coat or laser-cut large complex shapes of CVD diamond material opens up opportunities for the use of diamond in new applications where diamond materials were previously unavailable or cost prohibitive. This section is intended to provide examples of unique wear applications in which diamond coatings and freestanding diamond offer unique cost-performance advantages. Because the specific requirements for wear part applications vary widely, this section is organized in unique wear part case studies including both thin-film and thick-film product concepts. The selected case studies include diamond-coated mechanical seals (Sec. V.A), freestanding CVD diamond dressing tools (Sec. V.B), and freestanding CVD diamond wire dies (Sec. V.C). A.
CVD Diamond-Coated Mechanical Seals
Cemented tungsten carbide materials such as WC-Ni are commonly used in a wide range of mechanical seal applications. Advanced ceramic materials such as aluminum oxide, silicon carbide, and silicon nitride are used in seal applications that can afford the trade-off in fracture toughness and strength relative to cemented carbide for the benefits offered by the unique properties
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of ceramic materials. Over the last decade, there have been ongoing efforts to develop CVD diamond coating on cemented tungsten carbide as well as nonoxide ceramic seals for selected applications. Because many of these efforts have been commercial in nature and performed under confidential terms within the industry, very little performance data has been released into the CVD diamond industry. Cline [25] has reported a friction study of diamond-coated mechanical seals that were engineered for a commercial aqueous pumping application. Although ceramic seal faces are often mated to a softer, lubricious, hardcarbon face in many applications, the study was focused on comparing the benefits of diamond-on-diamond versus SiC-on-SiC seal face combinations. This study was focused on unfinished, as-deposited, diamond-coated seals produced using four different CVD process conditions. The sintered ␣-SiC material was a commercially available grade commonly used in the seal industry.* All seal pairs were tested in a friction tester using light seal face pressure, a mean surface speed of 306 m/min, and a starting temperature of 25⬚C. Both lubricated (deionized water) and unlubricated (dry) operating conditions were investigated. It is interesting to note that the surface speed used is nearly the same as that used in the presintered ceramic machining test summarized in Sec. IV.B.2.c. The results of the friction study are summarized in Table 7. Because of the proprietary nature of the results, the friction data were averaged for the four unique, but similar, diamond coating types. The table shows the diamond-on-diamond seal face combination was clearly superior to the SiCon-SiC. The diamond coatings decreased the mean steady-state friction coefficient during both lubricated and dry operation. The friction coefficient measurements of the diamond-coated faces indicate that CVD diamond coating can be tailored to produce friction coefficient values comparable to those of well-known lubricious materials such as fluoropolymers.† Because of the added lubricity of the diamond-coated faces, the temperature rise observed during the worst-case scenario of dry operation was reduced by over 100⬚C to a temperature range that minimizes the risk of damage to other components within the aqueous pump. Abusive rig testing conducted after the friction testing resulted in minimal diamond erosion compared with that observed for the silicon carbide face combination. Extensive application testing over a period of years indicated that the unique combination of properties
*Carborundum’s Hexoloy SA sintered ␣-SiC material. † According to E. I. du Pont de Nemours and Company, the coefficient of friction of Teflon is generally in the range of 0.05 to 0.20, depending on the load, sliding speed, and particular Teflon coating used.
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Table 7
Dynamic Friction Test Results of CVD Diamond-Coated Mechanical Seals
Seal face lubrication
Mean steady-state friction coefficient (ring-on-ring)
SiC-on-SiC assembly
Lubricated, deionized water
0.08–0.15
SiC-on-SiC assembly
Unlubricated, dry
0.20–0.30
Diamond-on-diamond assembly Diamond-on-diamond assembly
Lubricated, deionized water
0.03–0.04
Unlubricated, dry
0.05–0.13
Seal face material combination
Seal face temperature rise 6–8⬚C above bulk fluid temperature after 58 min of operation >350⬚C above ambient temperature after 5 min of dry operation 1⬚C above bulk fluid temperature after 58 min of operation 200–250⬚C above ambient temperature 14 min of dry operation
Source: Test data provided by a customer of Norton Diamond Film/Saint Gobain Industrial Ceramics, Inc.
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afforded by the CVD diamond coatings can eliminate several failure modes common to the silicon carbide seal faces commonly used in the application. B.
CVD Diamond Dressing Tools
Dressing tools are generally used to remove dull ceramic or superabrasive grains from grinding wheel surfaces in an effort to suppress grinding temperatures, improve workpiece finish, and improve workpiece tolerances. They are also used for ‘‘truing’’ or shaping grinding wheels. Historically, most dressing and truing has been performed with stationary dressing products that are fed into a rotating wheel. These products are often made using single-crystal or PCD diamond. Today, many high-productivity grinding operations employ rotary CNC dressers, or shaped dressing rolls, which can be used to dress one side of a grinding wheel continuously as the wheel is simultaneously grinding a workpiece on its opposite side. CNC rotary dressers may contain over 100 similar pieces of diamond that are hand-set over the outer diameter of the dresser and can be engineered to create unique forms in a grinding wheel. Examples of stationary and rotary diamond dressing tool designs are shown in Fig. 28. CVD diamond in a form similar to that shown in Fig. 29 offers several unique advantages to manufacturers of dressing tools as well as dressing tool end users. CVD diamond has been shown to offer performance enhancement for stationary dressing tools as well as rotary dressing tools in-
Figure 28 Examples of rotary (left) and stationary (right) diamond dressing tools used for shaping and truing grinding wheels. (From Ref. 1.)
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Figure 29 SEM images of a laser-cut piece of freestanding diamond typical of that used in a dressing tool. The 100⫻ image in (a) depicts an oblique view of the coarse-grained diamond facets common to as-deposited, freestanding CVD diamond. The 122⫻ image in (b) shows the relatively smooth nucleation face of the diamond (top center). The striations seen at the lower left of both (a) and (b) are the result of the laser-cutting process. The fractured cross section seen at the right end of (a) and front, center of (b) depicts the dense, polycrystalline structure of the diamond deposit. Note that the fractured cross section does not have the weak grain boundary characteristics (intergranular fracture) previously shown in Fig. 11a. (From Norton Diamond Film/Saint Gobain Industrial Ceramics, Inc. Images courtesy of Walter Swann; part (a) from Ref. 1.)
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cluding CNC profiling dressers and plunge form dressers. The commercial availability of synthetic, thick-film CVD diamond of good mechanical integrity creates a virtually unlimited supply of uniform laser-cut pieces that can offer superior performance relative to single-crystal diamond. Otherwise cost-prohibitive or unavailable, high-aspect-ratio pieces of diamond (such as 8 mm ⫻ 1 mm ⫻ 1 mm thick) can be cost-effectively incorporated into CVD diamond dressing tool designs allowing longer life and better dressing efficiency. Because of the precise shape of the laser-cut CVD diamond, the diamond can be strategically oriented to achieve 100% diamond consumption during the lifetime of the dresser. The availability of large-area CVD diamond helps to minimize or even eliminate design challenges associated with the crystal habit constraints and shape variability inherent in natural diamond. This enables small, intricate geometric forms such as sharp corners to be designed into dressing tools. The bulk structure of the CVD diamond can be engineered to enhance fracture toughness and strength for purposes of optimizing dressing tool life and dimensional stability beyond other forms of diamond (see Sec. III.A). This allows plunge form dressers to be produced with very small convex radii (e.g., 0.05 mm). C.
CVD Diamond Wire Dies
The production of metal wires is normally achieved by drawing metals such as copper, tungsten, or stainless steel through precision-made dies composed of hard materials such as diamond or tungsten carbide. Single-crystal diamond wire dies are much more wear resistant than tungsten carbide dies, but they have relatively poor chip resistance and are extremely difficult to fabricate. Furthermore, the intrinsic anisotrophy of a single crystal can result in nonuniform wear or cleavage-induced catastrophic failure under the extreme pressures of the wire drawing process. According to Wilks and Wilks [83], ‘‘the best choice of [crystallographic] direction is not too obvious because as the wire passes through the die its circumference is abrading the diamond on a whole 360⬚ range of planes, and the rates of wear on these planes will be somewhat different. Hence, the originally circular hole will not only grow larger but will loose its shape. However, 具110典 directions offer the advantage that the wire is abrading the sides of the hole with {001} and {011} orientations in abrasion resistant directions.’’ Over the years, efforts have been made to overcome the anisotropic trade-offs of single-crystal diamond wire dies by using polycrystalline diamond compacts [84,85], but these approaches were limited by the two-phase construction of a diamond composite and a secondary material of inferior mechanical properties. With the advent of CVD diamond processes, the ability to produce pure, dense layers of diamond created new
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alternatives for the wire drawing industry. Patents by Anthony et al. [86– 89] report the advantages of CVD diamond that has been microstructurally engineered for enhanced performance in wire drawing. Included in the claims of these patents is a 具110典 preferred, columnar orientation of the microstructure within the freestanding CVD diamond die. The polycrystalline diamond component of the wire die is oriented in a manner crystallographically similar to that described by Wilks and Wilks [83] for singlecrystal wire dies where the {001} and {011} orientations are present in abrasion-resistant directions. Although specific performance advantages are not presented here, the potential utility of CVD diamond synthesized in an oriented manner is quite clear in this application. D.
Other Wear Part Opportunities
Although cutting tool performance information is frequently reported in trade journals, trade magazines, and marketing literature, it is not feasible to predict the true state of the technology for the full range of mechanical applications where CVD diamond offers unique solutions. This is due mostly to the nature of commercial competition in the market. The introduction of commercial products for many wear part applications is likely to accelerate as CVD diamond production costs are reduced and the industry matures. Many wear part applications such as those described here require attention to detail specific to the application. Although much of the activity in the CVD diamond marketplace presently appears to be focused on larger, less fragmented markets, it is likely that the use of CVD diamond wear products will evolve and or be exposed in time. This trend will be driven by a necessary link between (1) a more thorough understanding of the structureproperty-performance relationships of CVD diamond design concepts and (2) commercial processes capable of reproducibly fine-tuning the necessary characteristics of the specific application.
VI.
SUMMARY
For decades, diamond has played a key role in industrial productivity. The 1950s saw the introduction of synthetic diamond grit. In the 1970s, metalsintered polycrystalline diamond (PCD) gained a foothold, offering unique flat shapes previously not available. Today, the evolution of diamond for mechanical applications continues with the emergence of CVD diamond as a versatile industrial material for cutting tools and other wear components. In machining applications, CVD diamond’s superlative combination of properties offers not only the promise of longer tool life under standard operating
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conditions but also opportunities for cost-effective high-speed and reducedcoolant operation.
A.
Thick-Film CVD Diamond
The inherent purity of CVD diamond deposition processes and industrial advances in freestanding CVD diamond manufacturing provide outstanding machining solutions for many difficult-to-machine materials. In most applications, freestanding, thick-film, CVD diamond tips are used as an alternative to cobalt-sintered PCD or single-crystal diamond. The ability to ‘‘tailor’’ the fracture toughness, strength, and other critical properties of freestanding CVD diamond is paramount to the ongoing development and commercialization of high-performance diamond cutting tool and wear products. Cutting tool application results indicate that the bulk mechanical properties of thickfilm CVD diamond can meet the needs of applications ranging from sintered tungsten carbide turning to high-silicon aluminum milling. The more recently discovered utility of thick-film CVD diamond for enhancing rotary and stationary dressing tool performance is a clear indication that the mechanical integrity of commercial grades of CVD diamond is suitable for very aggressive operations. The application range for thick-film diamond will continue to expand as diamond cutting and dressing tool fabricators gain a greater practical understanding of the unique properties of CVD diamond relative to PCD, single-crystal diamond, and other superhard materials.
B.
Thin-Film CVD Diamond
As practical methods continue to be developed for harnessing CVD diamond’s outstanding properties in diamond-coated products, the application range of thin-film CVD diamond continues to expand. Now that CVD diamond reactor development is reaching a level that supports commercialization and mass production, the primary technical challenge in the manufacture of diamond-coated carbide cutting tools and wear parts is ensuring suitable, reproducible diamond adhesion. A broadening portfolio of industrial adhesion technologies has led to the expanded use of diamond-coated cemented carbide tools in a wide range of metals and especially nonmetals. Depending on the vendor, the carbide grades used for diamond coating may be stringently controlled to ensure coating quality, adhesion, and long cutting life. As adhesion treatments become more robust, the ‘‘hidden promise’’ of CVD diamond-coated cemented carbide will finally be universally realized in the marketplace.
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ACKNOWLEDGMENTS Both authors are previous employees of the Norton Diamond Film business unit of Saint Gobain Industrial Ceramics, Inc. in Northboro, MA. The authors would like to acknowledge those who directly and indirectly contributed to this chapter including Saint Gobain coworkers, past customers of Norton Diamond Film, as well as academic mentors at Penn State University, Worcester Polytechnic Institute, and Bates College. Erica Olson and Heather Cline are acknowledged for their patience and support of the respective authors during their academic and industrial endeavors in the CVD diamond field. Portions of this work were funded by a NIST/DOC ATP program entitled Accelerated Commercialization of Diamond-Coated Round Tools and Wear Parts, an R&D joint venture between Kennametal, Inc. and Norton Diamond Film. REFERENCES 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14. 15. 16. 17. 18. 19. 20. 21.
B Cline. Cutting Tool Eng 52:76, 2000. BL Cline. Metalwork Equipment News March:68, 1999. RB Aronson. Manuf Eng January:52, 1999. KJA Brookes. World Directory and Handbook of Hardmetals and Hard Materials. 5th ed. Hertfordshire, UK: International Carbide Data, 1992. PL Smith. Am Mach December:74, 1999. J Schneider. Manuf Eng January:66, 1999. PM Stephan. Ceram Bull 71:1623, 1992. RA Hay. The new diamond technology and its application in cutting tools. In: ED Whitney, ed. Ceramic Cutting Tools. Park Ridge, NJ: Noyes, 1994, p 305. DB Arnold, FJ Momper. Manuf Eng November:62, 1997. DT Quinto. Int J Refractory Metals Hard Mater 14:7, 1996. WR Pfouts. Manuf Eng 125:98, 2000. KE Spear. J Am Ceram Soc 72:171, 1989. WA Yarbrough, R Messier. Science 247:688, 1990. JP Dismukes, KE Spear, eds. Synthetic Diamond: Emerging CVD Science and Technology. New York: John Wiley & Sons, 1994. BL Cline. Metalworking Equipment News January:24, 1999. K Kanda, S Takehana, S Yoshida, R Watanabe, S Takano, H Ando, F Shimakura. Surf Coat Technol 73:115, 1995. SY Kazutaka Kanda, K Yoneshima, H Ando. Diamond Films Technol 8:93, 1998. D Myers. Moldmaking Technol October:35, 1999. D Myers. Moldmaking Technol November:42, 1999. D Myers. Moldmaking Technol December:33, 1999. RS Sussmann, JR Brandon, SE Coe, CSJ Pickles, CG Sweeney, A Wasenczuk, CJH Wort, CN Dodge. Finer Points 10:7, 1999.
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11 Diamond Heat Spreaders and Thermal Management Ajay P. Malshe and W. D. Brown University of Arkansas, Fayetteville, Arkansas
I.
INTRODUCTION
Almost from the instant that active electronic devices were invented, a drive to increase the operating speed of electronic systems containing them was initiated. At the same time, a parallel effort focused on packing more electronics into a smaller and smaller volume. These trends continue unabated today. However, it is becoming more and more difficult to make progress in these areas using conventional single-chip packaging and present-day thermal management technologies. Multichip module (MCM) packaging, in which integrated circuits (ICs) are mounted extremely close together on a single substrate in order to increase the operating speed of an electronic system, only contributes to the thermal management challenge. As a consequence of these continuing trends in electronic systems, there are three main difficulties that must be addressed: how to provide for (1) the required power at the proper locations in the system, (2) sufficient input/output (I/O) connections to and from the system, and (3) adequate heat removal from the system. A.
Need for Thermal Management in Electronic Packaging
In IC technology, increasing functionality and operating speed of a chip are accomplished by reducing the minimum feature size, thereby shrinking device size and interconnection (i.e., wire) width. The reduced device dimen511
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sions dictate a reduction in the operating voltage(s) in order that material stress levels not be exceeded, thereby preserving the reliability of the IC. A reduction in interconnect width is accompanied by an increase in interconnect resistance, which then must be alleviated by reducing the interconnect length. On an IC, reduced interconnect length is a direct consequence of reducing device size. The final result of these changes is that an IC contains more active devices in a given area operating at higher speeds. This, in turn, means that the power consumption/dissipation in the given area also increases as illustrated in Fig. 1. Thus, a higher functionality IC with its accompanying higher power dissipation presents a thermal management challenge. As the drive to pack more and more of these ICs into smaller and smaller volumes without adversely affecting the operating speed continues, thermal management of electronic systems will present formidable challenges. The multichip module, noted previously, is a technology presently being used to pack ICs into a smaller volume by mounting them on a single substrate. The substrate provides both mechanical support for and interconnections between the ICs. By using array I/O on the MCM substrate, both the large I/O requirement and power insertion problems are relieved. However, thermal management becomes an even bigger problem because, as the power dissipated in a given volume increases, conventional approaches to dealing with the resulting heat quickly become inadequate. For example, the size of conventional heat sinks would have to increase to handle higher
Figure 1
Trends in clock speed and power consumption. (From Ref. 95.)
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power chips. This, in turn, would lead to a requirement for higher cooling air velocity and flow. However, in order to maintain or even reduce the size of a given electronic system, larger conventional heat sinks are not acceptable because of the increased volume and weight that they present. On the other hand, if the excess heat is not dissipated in an effective and timely fashion, there is the danger of degrading the system, which, at the very least, would adversely affect the long-term reliability of the system or, at worst, would prove fatal. Because of the increasing demands on packaging of electronic systems and the accompanying thermal management challenges, in recent years much attention has been directed at the development, processing, and characterization of new electronic materials and thermal management technology. Thus, the requirements and role of heat spreaders (also referred to as thermal spreaders) and heat sinks and the various materials and schemes used for thermal management of electronic packages are topics of a significant amount of research and development. B.
Thermal Heat Spreaders and Heat Sinks
A thermal heat spreader material serves as an interface between a heatdissipating device, such as an IC, and a heat sink. The primary function of the heat spreader is to collect heat from the source (e.g., IC) and quickly spread it to a larger device (e.g., heat sink) in order to eliminate heat from an electronic system. Thus, in terms of the roles they play in removing heat from an electronic system, the function of a heat spreader is different from that of a heat sink. In other words, the role of the heat spreader is to extract the heat rapidly from the source and spread it over a large area; it is the heat sink that actually dissipates the heat. Because of their different roles, a heat spreader needs to have a high thermal conductivity and excellent electrical insulating properties and should be capable of withstanding the high temperatures generated by the heat source device. To ensure intimate contact between the heat source and the heat spreader, the surface of the heat spreader in contact with the heat source should be extremely flat, of the order of or exceeding 0.1 m/mm, and the material should be stress free. Stress generally causes surface curvature, which is undesirable. Because of the requirement for an extremely flat mating surface, inexpensive processes for surface finishing, such as polishing or planarization, should be available and procedures for dicing or cutting should be simple and clean. Also, the definition of complex metallization patterns using photolithography should be manufacturing transparent. Apart from these features, the cost of the heat spreader must be sufficiently low as to make it economically viable.
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The physical dimensions of a heat spreader required for a specific application can be estimated by modeling. Usually, the thermal properties of the material used for the heat spreader, the configuration of the system generating the heat, and the amount of heat generated by the system are the input parameters to such a simulation program. The results from one such program involving the estimation of isotherms in half-cylindrical heat spreaders of different thicknesses are shown in Fig. 2 [1]. Single-crystal, natural diamonds have excellent properties for heat spreader applications. They have a very high thermal conductivity of over 20 W/cm K and electrical resistivities in excess of 1013 ⍀-cm. However, natural diamonds have a number of limitations that restrict their use in electronic applications. They are scarce, making them expensive, and from a more technical point of view, their composition is likely to vary from sample to sample. Furthermore, their physical dimensions severely limit the area to which they can be applied and, hence, their useful applications. To overcome these limitations, diamond substrates can be synthesized by chemical vapor deposition (CVD). Although CVD diamond deposition rates are reasonably high, the costs associated with its growth and postdeposition processing are still too high to allow it to be used routinely as a thermal management substrate in either MCMs or single-chip packages. However, given that diamond will probably be a viable MCM substrate in the future, a primary technical consideration is the minimum diamond substrate thickness required for efficient thermal management. As seen in Fig. 3, the efficiency of diamond in spreading heat decreases with increasing thickness for an edgecooled package geometry. This fact has important technical as well as economic implications when determining the optimum thickness of a substrate. Currently, approximately 1000-m-thick substrates are considered acceptable for MCM applications. However, according to some simulation results [2], thinner diamond substrates (20 to 30% thinner) can be used to achieve similar thermal management performance. This finding suggests that the CVD growth time and correspondingly the cost of diamond can be significantly reduced. Thus, CVD diamond, by far the technically best material for heat spreader applications, has the potential to become a viable contender economically. C.
Materials and Schemes for Thermal Management of Electronic Packages
Of the several conventional MCM substrates used for thermal management, silver (thermal conductivity, k = 4.28 W/cm K), copper (k = 3.96 W/cm K), beryllium oxide (k = 2.23 W/cm K), and aluminum nitride (k = 2.30 W/cm K) are the only materials having thermal conductivities close to that of
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Figure 2 Isotherms in half-cylindrical heat spreaders of different thicknesses. (From Ref. 1.)
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Thermal management efficiencies of some materials. (From Ref. 2.)
diamond. Figure 3 compares the thermal management efficiencies of these materials and CVD diamond, with k values as high as 16.0 W/cm K, although it can vary over a wide range of values depending on the growth conditions and it is anisotropic with the higher value in the direction parallel to the growth direction. It shows that even under extreme power dissipation conditions, the temperature rise at the center of a diamond substrate of optimum thickness in an edge-cooled geometry is much less than that of conventional materials. Moreover, the temperature rise for diamond is well within the acceptable limits for all practical applications (see Fig. 4). Many different schemes for thermal management of electronic packages are possible, depending on the particular application requirements. These schemes can be broadly classified into two categories: cooling of the package from the back and cooling from the edges. Although diamond has the highest thermal conductivity of any material, it has a relatively low heat capacity. Therefore, to use diamond to its best advantage, it should be used as a heat spreader by incorporating it as an intermediate layer between a heat source and a heat sink. Thus, as noted previously, the primary function of the diamond would be to spread the heat quickly from a smaller area (the heat source) to a larger area (the heat sink) where the heat can be disposed of more effectively and efficiently (see Fig. 5). In an application as a heatspreading substrate, the heat source (e.g., an IC) would be mounted on one
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Figure 4 Temperature distribution of plot of diamond substrate—simulation results. (From Ref. 2.)
Figure 5 Ref. 4.)
Placement of diamond film between the source and sink of heat. (From
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flat surface of the diamond with the heat sink mounted on the other side. In this configuration, the CVD diamond would quickly pull the heat from the IC, spread it over a large area (the diamond heat spreader), and transfer it to the heat sink, which is more efficient than diamond at dissipating heat. The fact that diamond has an excellent thermal conductivity but relatively low heat capacity reinforces the need to determine the optimum thickness for a particular application, thereby minimizing the economic impact of using a diamond heat spreader. Maximum performance benefits of diamond heat spreaders occur when the lowest junction operating temperature of the device generating the heat is achieved for a given configuration. On the other hand, minimum cost is realized by making the diamond substrate as small as possible because the cost of CVD diamond correlates directly with its volume. It has been observed [3] that the optimum diamond heat spreader should have a thickness that is 0.5 to 1.0 times the radius of the heat source and a radius three times that of the heat source (see Fig. 5). It has also been found that [3] when a CVD diamond–coated substrate is used as a heat spreader, the coating is effective only when the heat source radius is comparable to the diamond thickness. Unfortunately, the layer of solder used to attach the heat source (e.g., IC) to the diamond coating can present a major thermal impedance, thereby reducing the effectiveness of the diamond in reducing the temperature of IC. Another feasible approach to packaging electronic devices, such as solid-state lasers, is by stacking them vertically to effect a two-dimensional array instead of the conventional approach where devices are spread out in the horizontal direction over a single heat spreader creating a simple onedimensional array. The only efficient method for removing heat from the two-dimensional configuration is horizontally through heat spreaders placed between the devices. The thermal management problem is then one of heat flow in one dimension along a horizontal bar. Similarly, three-dimensional MCMs, which offer the greatest potential for speed performance enhancement of electronic systems, can be realized by taking advantage of the excellent thermal properties of CVD diamond. However, they also present the greatest challenges to the package designer and substrate fabricator. D.
Diamond Heat Spreaders
Effective, economical thermal management is essential for satisfactory operation of both single-chip and multichip electronic packages. Diamond has an excellent thermal conductivity (k = 20.00 W/cm K for natural diamond and in the range of 7.00–16.00 W/cm K for thermal management quality CVD diamond), a low coefficient of thermal expansion (⬇10⫺6), high elec-
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trical resistivity (>1013 ⍀-cm), and high mechanical strength (Young’s modulus ⬇ 1011 N m⫺2), making it an ideal choice for two-dimensional as well as three-dimensional MCM packaging. Thus, diamond is clearly a qualitatively superior material, compared with its competitor materials such as silicon, alumina, and beryllium oxide, for thermal management applications. In fact, it has been shown to perform well as a substrate for specific devices, such as high-power semiconductor laser diodes, and has been used as the substrate in prototype MCMs. The only obstacles to more widespread use of diamond substrates in such applications have been its limited availability in large sizes, which are stress free and of uniform thickness, and cost. However, advances in CVD diamond deposition technology make it possible to fabricate fairly large area (10 cm2) diamond substrates in the laboratory at a reasonable cost, thereby increasing its market considerably (see Fig. 6). The superiority of diamond as a thermal management material arises primarily from its high thermal conductivity combined with its electrical insulation properties. The purest natural diamonds (type IIA) have a thermal conductivity of about 20 W/cm K at room temperature. Those containing nitrogen as an impurity (type I) have a lower thermal conductivity. CVD diamond shows a variation in conductivity from near that of copper (4 W/ cm K) to about 16 W/cm K. Variations in the thermal conductivity are also observed when measured along and across the plane of the substrate. This variation arises from variations in crystalline size, chemistry, and crystal orientation. Although the thermal conductivity is slightly more in the direc-
Figure 6
Large-area (10 cm2) diamond substrates. (Courtesy of Norton Diamond.)
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tion perpendicular to the substrate surface, the in-plane thermal conductivity is still excellent in high-quality material (see Table 1), making diamond an excellent choice for thermal management applications. Furthermore, the low coefficient of thermal expansion is well matched to that of silicon, ensuring that stress is minimal when used with silicon ICs. Diamond is also a lowloss dielectric material, which is an important property for substrates used in high-speed digital circuitry. CVD diamond films or substrates are already competitive in price for applications requiring ‘‘small’’ pieces. For example, a small-area diamond submodule could be used as a drop-in module in a large electronic module. However, there are a number of applications where thin, large-area CVD diamond substrates could be used. Numerical calculations and simulations have shown that depending on the application, there is an optimum thickness for the CVD diamond film that will efficiently distribute the heat. This makes it possible to realize some devices and configurations that have not been possible to date. This includes multiple heat sources on a single substrate. Each heat source may be quite small, and provided the separation of the heat sources is sufficiently large, the required thickness of the substrate will be determined by the size (and power dissipation) of the individual heat-
Table 1
Measured Properties of CVD Diamond
Thermal conductivity (at 25⬚C) In-plane Perpendicular Optical transmission = 10 m = 2 m Young’s modulus Shear modulus Bulk modulus Poisson’s ratio Thermal expansion (RT 400⬚C) Density Specific heat at RT Hydrogen content Dielectric constant Loss tangent Electrical resistivity Flexural strength Chemical resistance Source: Ref. 4.
13.50 to 16.10 W/cm K 21.70 W/cm K >70% 63% 1180 GPa 514 GPa 558 GPa 0.148 2.0–2.6 ⫻ 10⫺6/K 3.51 g/cm3 0.54 J/g K 20 ppm 5.6 1011 ⍀ cm ⬃2 ⫻ 10⫺6/⬚C (25–200⬚C) 5.7 > heats the top surface of the specimen. The large beam size allows 1-D anal-
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Figure 12 Schematic diagram of the photothermal radiometry method for the measurement of . (From Ref. 94.)
ysis. A CaF2 lens focuses the IR radiation emitted by the center of the heated area onto an InSb detector. The thermal signal is measured as a function of . In another approach, the detector scans the thermal signal as a function of r, which is the distance from the modulating heating beam. In this case, the thermal signal primarily depends on the parallel component of the diffusion coefficient, i.e., on D㛳. Usually, a nonlinear fitting procedure is required to obtain the thermal diffusivity from thermal radiometry measurements. Obviously, the 3-D analysis is more complex than the 1-D procedure. d. Photothermal Deflection (Mirage Effect). A schematic diagram of the test setup for this measurement is shown in Fig. 13. The specimen is heated by a modulated laser beam that is focused normally onto the surface. The heating causes a thermal wave to propagate in both the specimen and the air surrounding the specimen. The temperature gradient in the air close to the specimen causes a change in the refractive index of the air in the immediate vicinity of the specimen. A probe laser beam, which is made to scan the heated surface of the specimen, is deflected by the index gradient in the air. The deviation of the beam depends on the distance of closest approach x between the probe beam and the heating beam. The deviation angle is projected into two angles n and t, which represent the component of the deflection normal to the specimen and the component parallel to the specimen, respectively. These deflections are measured with a quadrant detector. By measuring n and t as a function of x for different values of and fitting the data to theoretical functions, the thermal diffusivity can be obtained.
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Figure 13 Schematic diagram for the photothermal deflection method (the Mirage effect). (From Ref. 94.)
e. Round-Robin Measurements. The National Institute of Standards and Technology (NIST) brought people together from U.S. industries for the purpose of discussing the need to standardize the procedures for thermally characterizing CVD diamond. To address the issue of measuring the thermal conductivity of CVD diamond accurately, a working group on standardization of thermal conductivity measurement of CVD diamond was established [44]. It was agreed that standardizing the measurement of thermal conductivity () of CVD diamond was important. However, because a standardized measurement procedure did not exist at that time, the group decided to conduct a round-robin thermal conductivity measurement experiment in which several laboratories agreed to perform measurements of on a given set of specimens [43]. A considerable variation in measured values of resulted from this effort, independent of whether the measurements were made at the same laboratory using different techniques or at different laboratories. The variations have been attributed to various causes [45]. In some instances, a single laboratory obtained different values at several locations on the same specimen, indicating substantial inhomogeneity over the surface of the specimen. Some measurement techniques, such as the photothermal deflection and transient thermal gradient methods, have depth selectivity so that parameters measured on one surface differ from those measured on another surface because of differences in growth conditions at the two surfaces. Because diamond synthesis progresses through columnar growth, the density of voids in a sample is also a factor that affects thermal conductivity. Consequently, if the probed depth varies from one technique to another, then
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the measured parameter values affecting thermal conductivity () will be averaged over that depth. As the void density is expected to vary with depth, it is natural that the measured values would also vary if the measurement technique has depth sensitivity.
II.
POSTDEPOSITION PROCESSING OF DIAMOND HEAT SPREADERS
As noted previously, depending on the deposition parameters, CVD diamond films of different orientations, sizes, and thicknesses can be grown using a variety of techniques. However, the particular parameters used to grow diamond films dictate, to some degree, their thermal characteristics. In fact, diamond films are usually also nonuniform in thickness and grain size. Furthermore, they are composed of randomly oriented crystallites and exhibit intercrystal stresses unless grown in a very controlled manner [46]. A.
Postdeposition Processing
The nonuniform thickness of CVD diamond films, along with their surface roughness, can severely limit their thermal management efficiency due to poor contact between a heat source and the diamond and between the diamond and a heat sink. Photolithographic definition can also be limited by these film parameters. Furthermore, metallization adhesion can be adversely affected. To overcome these potentially serious problems, CVD diamond films must be subjected to postdeposition processing. The exact postdeposition processing steps required depend on the specific application of the diamond. For thermal spreader applications, it is essential that both surfaces of the diamond are flat and clean. B.
Cutting and Drilling
The typical freestanding diamond substrate produced commercially using state-of-the-art CVD techniques has a collar at its outer edge resulting from diamond growth extending down the side of the mandrill on which the diamond is grown [47]. Also, the outer portion of an as-grown substrate generally has a nonuniform grain size/chemistry if it is a large substrate. Consequently, the first required postdeposition processing is the removal of the collar so that the resulting substrate is planar. Furthermore, there is a need to cut the as-grown substrate into various sizes and shapes, as dictated by specific thermal management applications.
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Cutting of polycrystalline CVD diamond poses some challenges. The hardness of diamond, along with its chemical inertness, mechanical brittleness, and polycrystalline nature, presents serious challenges for precise cutting using conventional methods. The problems are compounded if the diamond film to be cut is thin and is deposited on another substrate that has different thermal properties. Typically, CVD diamond is cut using a pulsed neodymium yttrium-aluminum-garnet (Nd-YAG) laser, assisted by a reactive gas jet. The wavelength, power, and repetition rate of the laser, as well as the chemical quality and thickness of the diamond, affect the definition of the cut, the cutting-induced burr, contamination, melting, and microcracking of the material. In fact, the definition of a cut corner is a function of the laser wavelength. Laser cutting speed can be enhanced by having the right photochemistry at the diamond surface and by introducing the proper gas(es), such as a combination of a reactive gas, like oxygen, and an inert gas such as argon. Although a typical application of diamond in thermal management has the IC mounted directly on one of its sides, the same and/or other side might be used for high-density interconnection. Also, the edges of the substrate may be used to extract heat, depending on the cooling scheme employed. Furthermore, drilling of vias, which are filled with metal to provide electrical interconnection from one substrate surface to the other, may be necessary for some thermal management diamond substrates and requires that the diamond chemistry of the via surface be retained following drilling. These vias are typically 100–150 m in diameter, although other sizes are possible, depending on the particular application. Depending on the size and design of a diamond-based MCM, the number of vias required could vary from a few hundred to a few thousand on one substrate. The quality of a via is determined by its edge angle, wall contamination, and the definition of the top surface hole wall junction. Consequently, the drilling of vias in diamond films using conventional techniques has not had much success. Although drilling of diamond using lasers is not new, very little literature is available on the subject because of its potential commercial value. At the present time, an Nd-YAG laser operating at 1.06 or 53.2 m is used in the pulsed mode for drilling via holes in CVD diamond [47]. However, laser drilling of thousands of vias in a large-area, freestanding CVD diamond substrate can cause bowing of the substrate during drilling, thus making the process difficult. However, successful laser via drilling has been accomplished by drilling from both sides of the substrate [48]. In another novel process, fine laser drilling of diamond has been demonstrated by laser treatment of diamond in a reactive liquid medium [49]. It was shown that liquid chemistry plays an important role in laser machining.
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Polishing and Planarization
State-of-the-art physical and chemical methods of polishing and planarization of CVD diamond can be broadly classified as (1) nonreactive contact methods, (2) reactive contact methods, and (3) noncontact methods. Each method offers some advantages but also has some disadvantages. 1.
Need for Polishing
The rough surface of diamond substrates is an innate limitation of state-ofthe-art diamond growth techniques [50]. Consequently, the surface must be polished before the substrate will perform effectively in thermal management or other applications. Polishing is typically achieved by microchipping or dissolution or both. Hence, the factors to be considered when selecting a polishing method for a given application include the as-deposited surface quality, the polish rate (and hence the total time required to attain a specified surface smoothness), the minimum achievable roughness, surface contamination, as well as the equipment and operating costs. Different methods used for polishing the diamond film are described next. 2.
Polishing Techniques
For CVD diamond, the term polishing is typically used to describe material removal using any process. However, the total polishing process actually consists of two very different operations: lapping (coarse material removal) and polishing (fine material removal). In the discussions that follow, unless specified otherwise, a distinction is not made between the two operations. a. Mechanical Polishing. The mechanical polishing of diamond proceeds by grinding the samples on a flat metal wheel or scaife, which is usually made of cast iron. An abrasive powder, consisting of material harder than that to be polished, and a binder, often olive oil, are spread on or impregnated into the wheel. For polishing diamond, diamond powder is used as the abrasive material. The cast iron wheel is rotated at a high speed (>2500 rpm). A load is applied to force the surface of the diamond against the rotating wheel [51]. A coarse powder (>1 m) is used in the initial stage of polishing. This process is called lapping. Figure 14 shows an example of a mechanically lapped diamond surface. The primary advantages of postdeposition processing of diamond using a process such as mechanical polishing are that the process is simple and scalable. Also, there is no requirement for substrate heating or for reactive gases. Furthermore, the cost of the equipment is low. Moreover, there are virtually no limitations on the size of diamond samples that can be polished. Finally, the chemical quality of the film is not drastically changed after the
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Example of a mechanically polished diamond surface. (From Ref. 95.)
polishing process. There are, however, some limitations that should be weighed against the advantages. The polishing time can be quite long, up to several days, and the applied pressure on the substrates can cause microcracking. The polishing rates depend on the quality of the diamond and orientation. b. Thermochemical Polishing. This technique is based on the atomic dissolution of carbon into a hot metal. The technique was introduced by Grodzinski in 1953 [52], and Yoshikawa has used it in more recent times [53]. Yoshikawa’s setup is schematically depicted in Fig. 15. CVD diamond substrates are polished using a hot plate made of iron [53], nickel [54], manganese [55], or molybdenum [54]. The polishing rate depends on the diffusion of carbon atoms from the diamond surface into the hot metal plate. In general, it has been found that of these materials, iron yields the highest polishing rate [56]. However, as polishing progresses, the carbon content in the iron increases, which decreases the polishing rate. The plate temperature is required to be maintained in the range of 700 to 950⬚C. At lower temperatures, the polishing rate is low because of insufficient chemical reaction. At higher temperatures, some voids have been observed in the diamond surface being polished. Their number has been found to increase with temperature [54]. They are formed because of selective etching of diamond in the presence of oxygen. Thus, the polishing temperature must be optimized
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Figure 15 Ref. 95.)
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Schematic diagram of a thermochemical polishing apparatus. (From
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and the optimum temperature depends on the metal from which the plate is made. Thermochemical polishing is affected by the ambient. In vacuum, using iron, an etch rate of about 7 m/hr is the maximum possible. However, the finish is coarse. A smooth finish is possible in a hydrogen atmosphere, but the polishing rate is only about 0.5 m/hr. Consequently, it is common practice to start the polishing in vacuum to take advantage of the higher material removal rate and then finish the polishing in a hydrogen atmosphere to achieve a smooth finish. Figure 16 shows an example of a thermochemically polished diamond surface. Researchers at AT&T Bell Labs have used different metals, such as manganese, cerium, and, lanthanum instead of iron [57,58]. Under optimized conditions, etch rates as high as 50 m/hr have been realized using these materials. The cost of equipment required for thermochemical polishing is slightly higher than that for mechanical polishing. A vacuum chamber may be required. The operating cost depends primarily on the cost of the metal and the heating medium. The limitation on substrate size is similar to that for mechanical polishing. Particles from the metal plate that remain on the diamond surface are a major source of contamination. Also, due to the higher processing temperatures at the diamond-metal interface, the possibility of forming a thin nondiamond layer, such as graphite or diamond-like carbon, on the surface of the diamond exists. Finally, the processing time is of the order of several hours and the technique may not be suitable for thin films. c. Chemical-Assisted Mechanical Polishing and Planarization (CAMPP). Chemical-assisted mechanical polishing and planarization has
Figure 16
Thermochemically polished diamond surface. (From Ref. 95.)
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been shown to be a promising technique for fine polishing, not lapping, of diamond substrates [59]. CAMPP uses mechanical polishing in conjunction with chemicals to enhance the removal rate. Diamond under load comes in contact with an alumina base plate in the presence of oxidizing chemicals such as KOH or KNO. The diamond removal rate is directly proportional to the speed of the polishing wheel. Hard materials, such as diamond powder, may be added to the chemicals to enhance the polishing rate. The polishing rate using this technique is higher than that for mechanical polishing. Surface contamination of the substrate is less than that of either mechanical or thermochemical polishing. The technique, in principle, is also capable of polishing nonplanar surfaces. The polishing rate is a function of diamond chemistry. Figure 17 shows an example of a CAMPPpolished diamond surface. The required equipment and operating costs are only slightly higher than those for mechanical polishing. d. Laser Polishing. Contact polishing methods, such as abrasive polishing, apply stress to the diamond surface, which makes them impractical for polishing very thin (typically a few micrometers thick) diamond films or substrates. Moreover, most of the contact techniques are unsuitable for contoured surfaces or for selectively polishing small areas on larger surfaces. It is for such polishing needs that noncontact polishing techniques, such as laser polishing, prove beneficial.
Figure 17
Example of CAMP polished diamond surface. (From Ref. 95.)
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Laser polishing typically uses an Nd-YAG Q-switched pulsed laser [60], an excimer laser [61], or a combination of the two [62] to effect rapid polishing of diamond. The polishing is accomplished by irradiating the diamond surface with a pulsed laser operating at a repetition rate between 1 and 100 Hz. The polishing mechanism involves transient thermal oxidation and/or evaporative ablation of diamond [63]. A schematic diagram of a laser polishing system is shown in Fig. 18, and Fig. 19 shows the surface morphology of a diamond film following laser polishing. Laser polishing is well suited for coarse material removal over large areas [62]. Polishing proceeds by traversing the diamond film under a stationary laser beam. Traverse step size should be a fraction of the average grain size to ensure homogeneous irradiation of the diamond surface and hence a consistent surface topology after polishing. Depending on the laser beam wavelength and power, the laser pulses will heat the diamond film to a depth of a few submicrometers to a few micrometers. Depending on the polishing parameters, an approximately 150-nm-thick layer of nondiamond carbon may remain on the sample surface after polishing [64]. The polishing rate is limited by the grain size, ambient conditions, laser spot size, and laser power. It has been shown [65] that a high peak power (>1 GW/cm2) YAG laser in an oxygen ambient leads to a polishing rate of about 100 nm/min. The fastest coarse diamond polishing rate using a laser was demonstrated using multiple lasers on a freestanding thermal management substrate [62].
Figure 18
Schematic of laser polishing of diamond thin films. (From Ref. 95.)
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Figure 19 Surface morphology of a diamond film after being laser polished. (From Ref. 95.)
The time required to polish a diamond substrate depends on the size of the sample, the wavelength of the laser, the laser beam size, ambient conditions, and the scanning capability of the stage/laser. The surface finish attainable is limited by the quality of the different diamond facets and grain boundaries, which have different evaporation rates in response to the same laser radiation [66]. Moreover, if the polishing technique is applied to thinfilm diamond, the laser can generate shock waves, resulting in microcracks that can adversely affect the adhesion of the diamond to the substrate [67]. Laser polishing is probably the best method for rapid, coarse polishing of diamond. The technique does, however, have limitations for fine polishing. Although laser equipment costs are high, the technique is still economical because of its high polishing rates. Furthermore, as noted previously, laser polishing is attractive because it can be used on contoured surfaces. e. Ion Beam Polishing. Similar to laser polishing, ion beam processing is a noncontact technique that uses a focused ion beam. The polishing mechanism proceeds by sputtering of carbon atoms by the ion beam from the diamond surface reactively and/or nonreactively depending on the
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ion being used. A schematic of an ion beam polishing system is shown in Fig. 20. The diamond polishing rate depends on the orientation of the sample with respect to the angle of incidence of the bombarding ion beam. Diamond samples are typically heated to a temperature of about 700⬚C and bombarded with argon or oxygen ions. At normal incidence, the yield for 500 eV O2 ions is about seven times larger than the yield for 500 eV Ar ions. Oxygen ions oxidize the carbon, which possibly accounts for the higher rate of material removal. An etching rate of 20 nm/sec has been reported [68] in an oxygen atmosphere. The equipment required for ion beam polishing is very expensive. Although the sputtering rate depends on the energy of the ion beam, higher
Figure 20
Schematic of ion beam polishing setup. (From Ref. 95.)
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beam energies have a higher potential for producing surface damage and may lead to surface graphitization if the beam energy is very high. However, although the technique is expensive, it is clean and does not leave any chemical contamination on the diamond. The sample size is limited only by the size of the ion beam chamber. f. Reactive Ion Etching. Reactive ion etching (RIE) of diamond films has been reported [69] using O2 and H2 gases. The etching was performed in a DC plasma at an rf power of 200–300 W and at a pressure of ˚ /min have been 65–80 mtorr. Diamond etch rates on the order of 300–400 A achieved using 400 eV oxygen ions. Hydrogen ions are known to produce lower etch rates than oxygen ions. However, using oxygen, RIE etches at a rate about three times higher than ion beam polishing. RIE can produce surface contamination due to plasma heating and may cause anisotropic etching of the diamond. The cost of the equipment is relatively high and the specimen size is limited by the active plasma area. 3.
Planarization
As noted previously, in state-of-the-art CVD diamond synthesis, depending on the growth temperature and pressure conditions, favored crystal orientations dominate the competitive growth process. As a result, the grown films are columnar and have a polycrystalline structure and a rough surface morphology, as seen in Fig. 21. They also contain a number of microactivities, which are a consequence of their polycrystalline and columnar structure. These microcavities manifest themselves as surface pits when diamond is polished. Surface roughness and surface pits are undesirable features, predominantly because of the photolithography and via processing limitations they pose during fabrication of a thermal management substrate (as shown in Fig. 7 earlier). Generally, the surface roughness increases with increasing the film thickness. For 500- to 900-m-thick diamond substrates, the large surface roughness (tens of m or larger) limits the thermal management efficiency as well as the photolithographic resolution. Chemical and/or mechanical polishing is used to reduce the surface roughness, but polishing does not eliminate the surface pits as seen in Fig. 22, which shows a scanning electron micrograph (SEM) that compares an unpolished diamond surface with a chemical-assisted mechanical polished and planarized (CAMPP) CVD diamond substrate surface. As microcavities are a natural result of state-of-the-art diamond deposition techniques, surface polishing will not eliminate them because, as polishing proceeds, some surface pits are polished away, but others are created as microcavities in the bulk approach the surface.
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Figure 21 Atomic force microscopy (AFM) image of an unpolished diamond substrate. (From Ref. 95.)
In a novel approach to achieving planarization of CVD diamond, Malshe et al. [70,71] developed a technique that involves filling the surface pits with a material, such as a polyimide, spin-on glass, metal paste, or electroplated metal, that is compatible with microelectronic processing technology. It was shown that commercially available polyimides are especially attractive for effectively planarizing a diamond surface. The ability of a material to planarize a diamond surface is determined by its viscosity, solid content, and flow and shrinkage during cure. If desired, the filler layer can then be polished back to expose the diamond surface, although this polishing step has been shown to be unnecessary from a thermal performance standpoint [71]. The materials used as a filler material for planarization were polyimide PI-2721, PI-2610, and PI-2611 from DuPont [72]. The first layer was PI-2610, which has a low viscosity and, hence, is effective in filling the microcavities. By using a low-viscosity material such as PI-2610, the potential difficulty of air being trapped inside the microcavities is resolved,
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Figure 22 Scanning electron micrograph (SEM) comparing an unpolished diamond surface with a chemical assisted mechanical polished and planarized (CAMPP) CVD diamond substrate. (From Ref. 95.)
as the low-viscosity material can easily flow into the high-aspect-ratio microcavities and fill them up. This layer is partially cured and the diamond substrate is then treated with a coat of PI-2611. PI-2611 is more viscous than PI-2610 and completely planarizes the diamond substrate. Curing the PI-2611, layer along with PI-2610, completes the planarizing process. The adhesion strength of polyimide on diamond is reported [50] to be close to 4000 psi. The polyimide films exhibited excellent adhesion of CVD diamond even after subjecting the planarized diamond samples to thermal shock testing, as specified in MIL-STD-883. The reduction in adhesion strength was only about 33%, so the adhesion strength remained well above acceptable limits for packaging applications. The surface roughness of polyimide planarized diamond is quite adequate for photolithographic processes. Moreover, it was shown by simulation studies [71] that a polyimide thickness of about 5 m does not significantly degrade the thermal performance of the diamond substrate.
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Metallization
Metals are used in the packaging of ICs to provide electrical interconnection paths for signals, power, and ground. The most popular metals used in packaging are gold, silver, copper, and aluminum. Copper readily oxides and cannot be deposited onto many polymers because it reacts with them. Aluminum is also easily oxidized, although the oxidation process is self-limiting. Gold, on the other hand, is a very attractive metal for applications as a primary conductor as it is one of the best electrical conductors known, has good thermal conductivity, is essentially inert in most ambients, is very bondable, and provides a low contact resistance. However, due to its high cost, gold is normally used only as a top layer metal. 1.
Requirements
Metallization of CVD diamond is essential for bonding ICs to the diamond and providing electrical interconnects between the ICs. Bonding of ICs to diamond using metal is necessary for electronic packaging that is intended to take advantage of the excellent thermal and electrical properties of diamond. A metal contact on diamond is generally required to be both strongly adherent and stable at high temperature. To obtain strong adhesion, unless the metal is a carbide former, it is necessary first to deposit a layer of metal (which acts as the adhesion layer) that chemically reacts with diamond. It is the stability and chemistry of this layer that essentially determine the hightemperature thermal stability of the metallization system. The interface layer generally must satisfy three requirements: (1) interdiffusion of the interface metal into the primary conductor and the substrate should be low, (2) the interface metal should be thermodynamically stable with respect to the primary conductor and the substrate, and (3) the interface metal should adhere well to both the primary conductor and the substrate and should be of uniform thickness. No single metal meets all these requirements for all substrates. Depending on the physical and chemical mechanisms involved, diffusion barriers may be classified into three categories: passive barriers, sacrificial barriers, and stuffed barriers. In a passive barrier, the interface metal is chemically inert with respect to the substrate and the primary conductor. In a sacrificial barrier, the interface barrier reacts with the substrate and/or the primary conductor in such a way that it will not be consumed by the reaction until the lifetime of the product has expired. For a stuffed barrier, the grain boundaries, which provide the easiest diffusion paths, are plugged. The top metal layer of a multilayer metallization system must provide (1) good electrical conductivity, (2) corrosion resistance, (3) strong adherence to dielectrics, and (4) compatibility with the assembly process. Gold is
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an excellent material for top layer metallization because of its thermodynamic stability as well as its compatibility with die attachment and wire bonding processes. Two schemes have been used for metallization of diamond substrates. In one of the approaches, the adhesion layer is chromium (Cr) or titanium (Ti), with the overlayer metal being either gold (Au) or a gold-tin (Au-Sn) alloy. This system will fail at high temperature because Cr and Ti are soluble in Au and, therefore, diffuse into the gold, resulting in degradation of the diamond-metal bond and loss of adhesion [72,73]. In another system, the adhesion layer and the overlayer are the same material, but they are separated by a diffusion barrier. The diffusion barrier layer is typically an amorphous film of Ti ⫹ Si ⫹ N that is reactively sputter deposited onto the adhesion layer before depositing the overlayer. The best results have been achieved using Ti for the adhesion layer and overlayer with a barrier layer of Ti ⫹ Si ⫹ N [74]. 2.
The Metallization Process
Metallization for electronic applications is challenging because it demands optimization and control of the process for consistency and repeatability. Furthermore, thermal management applications require thick metallization, which has to be economic. One of the most reliable and economic techniques for realizing reliable and consistent metallization is electroplating [75]. However, extra effort is required when electroplating an insulating material such as CVD diamond because insulators cannot be used as an electrode in an electroplating bath. Hence, an electrically conducting seed layer must first be deposited onto an insulator before it can be electroplated. Deposition of a seed layer is generally accomplished by sputtering or by thermal evaporation. Although a seed layer is required, electroplating is one of the best, most reliable, and most cost-effective methods for metallizing CVD diamond. Wang et al. [76] discuss a method for metallizing plated through holes (PTHs) using electrolytic copper. To ensure uniform plating inside PTHs that are 13 mils or smaller in diameter, they use a very high throwing power. This is realized by using a low copper concentration and a high acid concentration in the plating solution. The throwing power can be further enhanced by using a low current density during the electroplating. In this way, it is reported that through holes as small as 6 mils in diameter can be plated [76]. Baker [77] described plating baths for the alkaline gold sulfite system in which arsenic was used as an additive to produce electroplates with hardness depending on the amount of oxalic acid added to the electrolyte. Larger amounts of oxalic acid produced lower hardness in the electrodeposit.
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Thick Film Metallization
Dupont Electronics Materials (Research Triangle Park, NC) and Crystalline Materials Corporation (San Ramon, CA) have reported on their development of thick-film metallization for CVD diamond, specifically for MCM applications [94]. Thin-film metallization has traditionally been the approach to providing metallization on CVD diamond. This newly developed thick-film metallization system, which uses a unique inorganic binder, is designed for nontraditional substrates such as diamond [92]. 4.
Via Filling
The use of at least one side of a diamond substrate for direct bonding of ICs, the same and/or other side for high-density interconnections, and the edges for heat extraction is a realistic configuration of a single high-density substrate used in a 3-D multichip module. However, this type substrate structure requires the use of vias to provide paths for electrical interconnection between the two sides. The material used to fill the vias obviously needs to be a good conductor, should have a coefficient of thermal expansion close to that of diamond, and should also have excellent adhesion to diamond. Tungsten-copper composite via-fill material has a reasonably good conductivity and a coefficient of thermal expansion more closely matched to diamond than most other conductors [48]. Raytheon Company (Sudbury, MA) and Micro Substrates Corporation (Tempe, AZ) have reported the results of their combined efforts to realize good via-filled diamond substrates [50]. In order to minimize the diamond substrate polishing effort, the smoother, mandrel side of the diamond substrate was used to form the metal interconnects by photolithography. A via-filled diamond substrate is realized in the following sequence: (1) CVD diamond is synthesized on suitable mandrels, (2) the diamond is separated from the mandrel, (3) the diamond is cut and polished, (4) vias are drilled into the substrate, (5) an adhesion layer is deposited in the vias because of poor adhesion between the via wall and conventional via-fill material [50], and (6) the via-fill material is deposited. The via-filled diamond substrate is then processed further to reestablish the resistivity of the diamond substrate. The following sections provided more detail on the process steps required to form filled vias. a. Laser Drilling. An Nd-YAG laser, operating at the fundamental wavelength (1.06 m and 532 nm), is found to be suitable for drilling holes in diamond, a few hundreds of micrometers thick. The process is attractive because it is clean and good-quality circular holes can be drilled. For drilling holes in thicker CVD diamond substrates, a process that uses laser core drilling from both sides of the substrate results in clean, circular holes.
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A CO2 laser, which is often used for high-rate drilling of alumina substrates, is not very effective for drilling holes in CVD diamond substrates due to the very low energy absorption of the CO2 wavelength by diamond. However, better results are obtained with CO2 lasers in the presence of oxygen. b. Adherent Deposits. Adhesion of the via-fill material to diamond is generally so poor that via plugs can be pressed out of their holes very easily. To improve the adhesion, an adherent layer of titanium-tungsten (TiW) can be deposited between the diamond and the via-fill material. Other materials, some proprietary, have also been used as adhesion promoter materials. The adhesion material is sputtered into the vias. Limiting the adhesion promoter layer to the via wall is important because it eliminates the need to repolish the diamond surface after deposition of the via-fill material. An etch-back process, involving a simple via-fill masking process that does not use any photomasks but uses a high-viscosity filling material, has been developed [48]. The material fills the via to the surface of the opposite side. A two-step fill, etch, and strip sequence provides good results. Completion of the etch can be determined by observing conductivity using microprobing techniques. c. Metallizing Vias. For small vias with a large aspect ratio, lowpressure CVD (LPCVD) techniques [79] are used to fill vias ranging from 0.001 to 0.010 in. in diameter in 0.010- to 0.030-in.-thick CVD diamond (aspect ratios in the range 1–30) [80]. For large holes with low aspect ratios, a hybrid sputter-plating process is used. The critical issues for a via-filling process include adhesion and electrical conductivity. Of these, an adhesion problem can be handled by using adhesion-promoting deposits as described previously. For small holes with high aspect ratios, penetration into the via and porosity of the deposit are critical. Tungsten is generally chosen for via filling using LPCVD because it is a good carbide former and is the best electrical conductor among the transition elements (which are known to be very good carbide formers and hence are all potential candidates for metallization of CVD diamond). Copper and aluminum are also excellent choices for via filling but require Ti-W as an adhesion material. Heat treating a thick (>0.5 m) tungsten film on diamond to 1000⬚C results in severe adhesion degradation. The tungsten is almost completely delaminated leaving no trace of tungsten (W) or any of its carbide (W2C or WC) on the surface. In ˚ ) of tungsten, when heat treated at 960⬚C for contrast, a thin layer (⬃200 A 30–60 min, maintains its excellent adhesion to diamond. Following the deposition and proper heat treatment of a thin tungsten film, the diamond substrate can be cooled and thick layer of tungsten can be applied.
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d. Resistivity Restoration. This step is used to remove via-fill material that may have adhered to the substrate surface during the via-fill process. It is interesting to note that, although the diamond surfaces may exhibit increased electrical conductivity following a via-fill process, the fill materials are not responsible for the increase. The reason for the increased conductivity is the presence of graphite, which is formed when the diamond is elevated in temperature during the various via-filling process steps. Restoration of the diamond resistivity is achieved by placing the substrate in an oxygen plasma. The plasma oxidizes the graphitized diamond layers, thereby restoring the resistivity. However, care needs to be taken to ensure that the metallization in the filled vias and other locations on the diamond surfaces are not adversely affected. It should be noted that neither argon sputtering nor plasma etching is capable of restoring the preprocessing conductivity of the diamond.
E.
Die Attachment—Face-Up and Flip-Chip
In MCM technology, reliable die attachment implies that there is excellent electrical communication between the ICs and the substrate on which they are mounted. Two die attach technologies are presently being employed to attach ICs to thermal management diamond substrates. One is referred to as flip-chip die attachment and the other is the conventional face-up attachment using wire-bonding attachment for electrical connection between the ICs and the substrate. These two approaches are depicted schematically in Fig. 23.
Figure 23 Schematic of flip-chip and face-up (wire bonded) die attachment configurations. (From Ref. 82.)
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Flip-chip die attachment, which was originally developed by IBM in the 1960s and referred to as Controlled Collapse Chip Connect (C4) [81], is now finding more widespread use due to its distinct advantages of selfalignment, high assembly yield, and good reliability in small coefficient of thermal expansion (CTE) packages. The configuration also offers very high I/O capabilities and a minimum I/O signal path length for excellent electrical performance compared with wire bonding. However, flip-chip attachment is more expensive and it also suffers from problems such as conduction heat dissipation in substrate-cooled packages and structural fatigue failure in applications with large CTE differences. Although the flip-chip approach places the ICs closer to the substrate, the solder bumps present an overall bottleneck to heat flow because they do not present a continuous interface layer. Thus, the choice of a die attach technology must be guided by the package design required for a given application. 1.
Thermal Modeling Procedure
Three-dimensional finite element analysis (FEA) models have been developed that estimate the temperature buildup of a device. The inputs to such a model are material parameters of the major components of the electronic package system (i.e., die, substrate, and attachment layer). Brown et al. [82] included the solder bumps of a flip-chip when performing a flip-chip analysis because the localized conduction associated with each bump can result in hot spots in the die. Further, the solder bumps were constructed using rectangular elements but were given a thermal conductivity value calculated to account for the C4’s typical truncated spherical shape. The heat load was assumed to result from heat dissipated by ICs operating uniformly over the surface of the substrate. Cooling was modeled as a constant temperature plane on the bottom or the edge of the substrate, two commonly used configurations to connect a diamond heat spreader to a heat sink. Parameters of interest included the die material properties, die dimensions, solder material properties, solder bump diameter and shape, solder bump array pitch, underfill material properties, substrate dimensions, and substrate material properties. As a point of reference, a baseline case was established with the model parameters having values as listed in Table 3. The maximum temperature differentials between the die junction temperature (Tj) and the substrate cooling temperature (Tc) as well as the heat flow paths were noted from the finite element solutions to compare package performance as a function of the selected parameters. a. Die Material. The choice of die material (Si or GaAs) does not have a large effect on either the face-up or flip-chip package thermal performance. In the face-up, bottom-cooled model, the normal resistance
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Base Case Model Parameters
Die material Die size Die thickness Die power Attachment material Face-up attach thickness Flip-chip bump height Solder bump conductivity Underfill conductivity Solder bump pitch Substrate materials
Substrate size Substrate thickness Cooling plane temperature
GaAs (45 W/m K) 10 mm ⫻ 10 mm 404.6 m 40 W (uniform) 95 Pb/Sn solder (35 W/m K) 25.4 m 101.6 m 44 W/m K (150 m diameter) 2 W/m K (high k filler) 357 m Diamond (1200 W/m K) Cu (400 W/m K) Si (148 W/m K) 15 mm ⫻ 15 mm 404.6 m 50⬚C
Source: Ref. 82.
through the die is kept small by the short conduction path and the edgecooled package resistance is dominated by the substrate. Although heat is not forced through the die in a flip-chip package, it is known that heat would flow through the higher thermal conductivity solder bumps to the substrate rather than through the underfill. The largest temperature difference was observed for an edge-cooled substrate, which resulted in the Tj of a silicon die being 13⬚C cooler than that of a GaAs die. b. Die Width. Simulation studies utilizing dies of different widths showed that while a bottom-cooled package saw a constant thermal conduction path and a proportional increase in heating and cooling loads, the edge-cooled package experienced an increasing thermal conduction path and only a linear increase in cooling area. Thus, the bottom-cooled package performance remained almost constant over a range of die widths, whereas the edge-cooled package experienced significant increases in Tj ⫺ Tc , which became more pronounced as substrate conductivity decreased. Figure 24 shows the flip-chip model results, which are very similar to the results for the face-up bonded case. If Tj = 100⬚C is defined as the die temperature at which device failure occurs, then a diamond substrate would be the only edge-cooled package that would survive over the full range of die widths considered.
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Figure 24 Maximum package temperature differential versus die width for flipchip die attachment. E-C, edge-cooled; B-C, bottom cooled. (From Ref. 82.)
c. Die Thickness. The effect of die thickness variation on the overall package performance is negligible when the conduction paths are small. Increasing the die thickness does increase the normal die resistance, but it simultaneously reduces the lateral resistance. Thus, in the face-up bonding case, the rise in the temperatures (Tj ⫺ Tc) is found to increase with increasing die thickness. In the flip-chip configuration, because it is the lateral resistance that is critical, cooler junction temperatures are achieved with increasing the die thickness. d. Die Power. Over the heat dissipation range of 20 to 80 W, which represents power dissipation levels in present-day and near-future highspeed, high-power electronics, all packages exhibit a linear increase in (Tj ⫺ Tc). Two approaches are commonly used for the cooling schemes: the bottom-cooled and the edge-cooled approach. Despite the temperature rise, a bottom-cooled substrate package remains at safe levels due to the short thermal paths. The edge-cooled diamond substrate is found to handle about 400% increase in power dissipation as compared with a Si substrate [84]. e. Flip-Chip and Face-Up Die Attachment Comparison. Flip-chip die attachment has a thermal advantage over face-up die attachment because of the presence of two paths for heat conduction: through the substrate and through the top of the die. Simulation results [83] show that, as the heat sink resistance decreases, the (Tj ⫺ Tc) for an edge-cooled substrate decreases toward the values for a bottom-cooled substrate. A die heat sink is
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essential for adequate cooling of an IC if a package uses a high-thermalresistance substrate without bottom cooling. However, using a diamond substrate eliminates the necessity for a die heat sink because the junction temperature would be maintained within safe operating limits. In other words, diamond is an extremely attractive substrate material for electronic packaging applications for thermal management of high-power-density ICs packaged into very small spaces.
III.
APPLICATIONS AND EXAMPLES OF VARIOUS PACKAGES
This section describes some of the passive electronic applications of CVD diamond as a thermal spreader reported in the published literature. A significant amount of progress has been made in the utilization of diamond in electronic packaging over the past few years, but challenges still remain. A.
Background
The U.S. Defense Advanced Research Projects Agency (DARPA) should be recognized for its continued interest in and financial support of research and development of synthetic diamond for electronic thermal management. There is no doubt that it would have been difficult, if not impossible, for many researchers to pursue diamond thermal management research at any level were it not for DARPA support. The following sections address the applications of diamond to electronic packaging. B.
Optoelectronics
Diamond finds applications in numerous optoelectronic applications. Conventional light-emitting diodes (LEDs) have a very high thermal resistance and a low light output. This makes them impractical for high-flux applications such as automotive tail lights. Work has been carried out at the Optoelectronics Division of the Hewlett Packard Company (San Jose, CA) in which LEDs were soldered to a diamond heat spreader mounted on a copper slug. The result is an enhanced ability to drive a high-light-flux LED, a larger LED chip at higher powers [90], resulting in a fivefold increase in the total light output per element. Table 4 compares the characteristics of conventional and high-light-flux LEDs. The enhanced performance achieved by HP may allow LEDs to compete on a cost per lumen basis with incandescent light sources in many applications.
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Table 4 LEDs
Comparison of Performance of Conventional and High-Light-Flux
Device Conventional LED High-light-flux LED
Thermal resistance (⬚C/W)
Light output (lumen)
Maximum drive power (W/cm2)
LED chip area (cm2)
300
1.8
250
0.00064
>316
0.00250
4
10
Source: Ref. 90.
From a manufacturing standpoint, processing precut, and generally small, diamond heat spreaders such as those used with LEDs proves to be difficult and expensive. A possible alternative is the use of relatively large diamond wafers after they are partially etched to define the edges of specific heat spreader sizes. Such wafers could then be broken into individual heat spreaders. This approach is reported [84] to have nearly 100% yield, is cost effective, and is convenient for postdeposition processing. Furthermore, gold-plated diamond heat spreaders have been successfully attached to silver-plated copper slugs using a thermocompression bonding technique [84]. This eliminates the need for an adhesive layer, which only adds to the thermal resistance of an electronic package. Diode lasers have already become an attractive choice for laser applications because of their low cost and high efficiency. Consequently, efforts continue to be directed at making diode laser arrays on a commercially viable scale that are reliable and cost effective. However, laser arrays result in an even greater need for efficient thermal management materials and cooling technologies. Not surprisingly, CVD diamond substrates are presently being used for thermal management of laser diode arrays. Sumitomo Electric USA Inc. has developed a composite diamond heat spreader material for laser diode applications [91]. C.
Three-Dimensional High Density Packaging
The multichip module approach to achieving very high packaging and interconnect densities in electronic packaging is of particular interest for digital system engineers. However, as the density of high-performance ICs in an MCM package increases, the power dissipation also increases, making thermal management a challenging task. Assuming that the cost of diamond can be reduced, it offers a promising solution to many difficult electronic
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packaging challenges, as can be seen from the comparison of the properties of electronic packaging materials in Table 5 [84]. Figure 25 illustrates the ability of diamond to quickly transport heat away from a heat source, such as an IC. 1.
Diamond for 3-D MCM Applications
There are several advantages in going to three-dimensional electronic packaging, but the most important one is a significant reduction in signal path length between contiguous 2-D multichip modules, thereby increasing the operating speed of a system. In 3-D MCM packaging, use is made of the fact that IC chips are thin (typically 0.6 mm or less), permitting very high volume density packaging. Figure 26 illustrates that in a 2-D MCM there can be only eight nearest neighbor ICs for a given IC. However, in a 3-D MCM, the same IC would have 116 nearest neighbors within the same distance. Assuming that the MCMs are 1.5 cm square and continuing the example further, if the interconnect delay is 150 psec/in. and the VLSI chips used have 10,000 gates, then the 3-D geometry places over a million gates within a worst-case delay time of about 650 psec. If these ICs were spread out horizontally, the worst-case delay would be in excess of 1.7 nsec, which
Table 5
Comparison of Electronic Packaging Material Properties
Material Diamond (natural) Diamond (CVD) Beryllia, BeO AlN Alumina, 99% GaAs Silicon Kovar (FeCoNi) Molybdenum Aluminum Copper Silver Diamond-epoxy Silver-epoxy Polyimide Source: Ref. 84.
Thermal conductivity (W/m⬚C)
Electric insulator? (Y/N)
Thermal expansion (ppm/⬚C)
2000 700–1700 220 70–230 29 45 149 17 146 237 396 427 8.7 5.8 0.2
Y Y Y Y Y Semi N N N N N N Y N Y
0.8–1.0 1.0–1.5 6.4 3.3 6.3 5.9 2.6 5.9 5.1 23.8 16.8 19.6 120 120 >50
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Figure 25 Ref. 84.)
Diamond capabilities in reducing ‘‘hot spot’’ temperature rise. (From
Figure 26 Ref. 84.)
Comparison of 2-D and 3-D MCM signal interconnect lengths. (From
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would represent a significant performance compromise for system clock rates in the 50–100 MHz range and virtually preclude operation at rates above 250–350 MHz. With 3-D packaging, this is extended to nearly 1 GHz. Diamond is attractive for 3-D MCM applications, consisting of vertically stacked 2-D multiple chip modules, primarily because of its extremely high thermal conductivity, a requirement for any substrate utilized in 3-D packaging of high-power ICs. Furthermore, technologies exist that will allow electrical interconnections to be made in the vertical direction (i.e., between the 2-D multichip modules). As shown in Fig. 27, conductive fuzz buttons can be used in spacer boards placed between the 2-D multichip modules to provide vertical interconnection. Such fuzz buttons are typically 0.020 in. in diameter and extend beyond both sides of the board by about 40 mils, thereby making solid contact with pads on the substrates above and below when a mating force is applied. Such an arrangement facilitates rework of defective modules. 2.
Edge Cooling
One possible approach to dealing with very high power density and the associated heat dissipation is to extend the diamond substrate and make it a part of the heat exchanger itself. As an example, if an active 100 mm ⫻ 100 mm MCM substrate is extended on the heat sink ends by 5 mm or so,
Figure 27
Cross section of 3-D diamond MCM assembly. (From Ref. 84.)
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then it is possible to perforate these extended ends with slots of 0.1 mm width and 2.5 mm length on a pitch of 0.25 mm to create a total surface area of about 50 cm2. This would reduce the power density from 250 W/ cm2 to about 50 W/cm2. This would make thermal management more effective and maintain the interface temperature at an acceptable level (see Fig. 28). The edge-cooling approach has been attempted experimentally in a joint effort by Sandia National Laboratories (Albuquerque, NM) and Allied Signal Corp. (Kansas City, MO) [87]. They have built and tested a single substrate intended for a 3-D edge-cooled MCM. Their results indicate that it is possible to achieve a thermal resistance as low as 0.22⬚C/W for a single substrate. The challenge is to achieve the same kind of thermal impedance for 3-D stacks of these substrates. D.
Power Packages
In recent years, there has been a trend toward higher power ICs while, at the same time, reducing their physical size. Typical power densities are now
Figure 28 84.)
Edge cooling technique utilizing extended cooling slots. (From Ref.
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on the order of 200 W/cm2 and rising. Power electronics poses one of the most difficult problems for miniaturization, but it should be possible even considering the challenges of thermal management by taking advantage of multichip module technology called multichip power modules (MCPMs). Although the components themselves are relatively small, often requiring only a few electrical connections, the package size is often determined by thermal loads and the desired thermal performance of the complete system. Diamond is an ideal thermal management material for such applications. However, if a layer of diamond is placed between a heat-generating device and a copper heat sink, the thermal shear stress acting at the diamond-copper interface far exceeds the shear strength of diamond. To overcome this problem, a 200-m-thick aluminum (Al, 99.995%) layer has been placed between the diamond and copper and also between the diamond and device, as shown in Fig. 29 [88]. Pure Al is very soft compared with copper and exhibits a plastic stress-strain behavior. The aluminum layer is thus able to absorb the thermal stress in the module. Figure 30 shows the temperature difference between the coolant and device surface for a module with aluminum oxide and diamond [89], clearly depicting the improvement in thermal management by using CVD diamond substrates. Diamond substrates are being used by Sandia National Laboratories (Livermore, CA; Albuquerque, NM) for MCM packaging of power electronic subsystems of flight computers [89]. The flight simulator computer is built around a radiation-hardened microprocessor, nonvolatile memory, and analog driver circuits and has multiple responsibilities including (1) providing a digital communication interface to the airplane or missile, (2) performing signal processing as required, (3) systems status monitoring, and (4) actuation signals to other systems (see Fig. 31). The actuation signals are large-power, short-duration signals for activation of mechanical and pyro-
Figure 29 Schematic cross section of a power module with diamond layer for thermal management. (From Ref. 86.)
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Figure 30 Photograph of working diamond MCM (dark substrate) and alumina MCM in copper tungsten secondary package. (From Ref. 87.)
technical devices. Two typical activation signals are a 150 W, 12 msec pulse and a 32 W, 20 Hz square wave of 3 sec duration. As the total activation time is brief, thermal performance is dominated by the first level of packaging around the active and passive devices. Temperature measurements made as a comparative study using diamond and aluminum nitride (AlN) as the thermal management materials revealed that in order to achieve the same temperature difference, the surface area of AlN required was about four times larger than that of CVD diamond substrates [89,90]. The effectiveness
Figure 31
Schematic of flight computer block diagram. (From Ref. 87.)
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of CVD diamond in thermal management applications involving a reduction of volume is thus obvious. The MCM version of the power electronics subsystem occupies a little more than half the volume of a conventionally packaged, through-hole PCB system. However, the shrinking of the power electronic subsystem does require the excellent thermal management capabilities of CVD diamond. In another approach, Vandersande et al. reported using thermoelectric coolers and diamond films for temperature control of power electronic circuits [93]. The idea is to mount the highest power components directly onto thin-film thermoelectric elements, which maintains the temperature of the components from a few degrees to a few tens of degrees below that of the diamond substrate. To maximize the efficiency of the thermoelectric cooler, diamond films acting as thermal lenses are used to spread the heat from the small power device to the larger coolers. However, thermoelectric coolers are not generally useful for temperature control of high-density electronic circuits. Passive cooling with high-thermal-conductivity CVD diamond films is a much more realistic method for cooling of electronic circuits above room temperature.
E.
Plastic Packages
Commercial, monolithic, microwave GaAs ICs (MMICs) are usually packaged in injection-molded plastic packages. When packaged in plastic, these devices are restricted to power dissipation level of 1–2 W. If the power dissipation exceeds these levels, more expensive ceramic packages are used. By integrating a diamond substrate into the plastic package, the thermal dissipation capabilities can be increased to about 15 W at about half the cost of a ceramic package. Thus, a combination of the superior material properties of CVD diamond and the economics of plastic packaging offers a potential solution for relatively inexpensive packaging of microwave ICs. Norton Diamond Film’s NorCool (NorCool is a trade mark of Norton Diamond Film) package combines the excellent thermal properties of CVD diamond with the low cost of plastic packages. The NorCool package replaces the die paddle of a typical copper lead frame with a diamond substrate. The diamond substrate is physically attached to all the leads, thereby providing tremendous improvement in thermal performance, yet keeping the leads electrically isolated from each other [92]. The major problem associated with the attachment of GaAs MMICs to diamond substrates could be the coefficient of thermal expansion mismatch, as the CTEs of GaAs and diamond, respectively, are 5.7 ⫻ ppm/⬚C and 1.4 ppm/⬚C.
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Concluding Remarks
Recent success in the synthesis, processing, and application of CVD diamond has demonstrated its potential as one of the most promising thermal management candidates for challenging electronic applications. Significant efforts have resulted in optimization and commercialization of synthesis and processing technologies. Emphasis should continue to make CVD diamond synthesis and processing more cost effective as well as manufacturing transparent. Diamond films and substrates will continue to find more and versatile applications for the ever growing thermal management needs for miniaturization, increased operating speed, and cost reduction in electronic packaging.
G.
List of U.S. Patents Relevant to Diamond Heat Spreaders 5,472,370 Method of planarizing polycrystalline diamonds, planarized polycrystalline diamonds and products made therefrom 5,663,595 Diamond heat sink comprising synthetic diamond film 5,660,894 Process for depositing diamond by chemical vapor deposition 5,653,952 Process for synthesizing diamond using combustion method 5,650,639 Integrated circuit with diamond insulator 5,648,148 Thermal management of electronic components using synthetic diamond 5,647,964 Diamond film synthesizing apparatus and method thereof using direct current glow discharge plasma-enhanced chemical vapor deposition 5,645,617 Composite polycrystalline diamond compact with improved impact and thermal stability 5,642,779 Heat sink and a process for the production of the same 5,633,088 Diamond film and solid particle composite structure and methods for fabricating same 5,631,046 Method of metallizing a diamond substrate without using a refractory metal 5,628,824 High-growth-rate homoepitaxial diamond film deposition at high temperatures by microwave plasma-assisted chemical vapor deposition 5,624,719 Process for synthesizing diamond in a vapor phase 5,620,512 Diamond film growth from fullerene precursors 5,614,258 Process of diamond growth from C.sub.70
Diamond Heat Spreaders and Thermal Management
571
5,609,955 Diamond film and mixed diamond and nondiamond particle composite compositions 5,609,683 Apparatus for making industrial diamond 5,607,723 Method for making continuous thin diamond film 5,607,560 Diamond crystal forming method 5,567,986 Heat sink 5,560,779 Apparatus for synthesizing diamond films utilizing an arc plasma 5,545,030 Diamond film and solid nondiamond particle composite compositions 5,539,176 Method and apparatus of synthesizing diamond in vapor phase 5,538,765 DC plasma jet CVD method for producing diamond 5,535,816 Heat sink 5,531,184 Method for producing synthetic diamond thin film, the thin film and device using it 5,523,160 Highly oriented diamond film 5,523,121 Smooth surface CVD diamond films and method for producing same 5,518,759 High-growth-rate plasma diamond deposition process and method of controlling same 5,516,554 Cyclic hot-filament CVD of diamond 5,514,242 Method of forming a heat-sinked electronic component 5,510,157 Method of producing diamond of controlled quality 5,508,230 Method for making a semiconductor device with diamond heat dissipation layer 5,507,987 Method of making a freestanding diamond film with reduced bowing 5,506,452 Power semiconductor component with pressure contact 5,505,158 Apparatus and method for achieving growth-etch deposition of diamond using a chopped oxygen-acetylene flame 5,504,303 Laser finishing and measurement of diamond surface roughness 5,500,157 Method of shaping polycrystalline diamond 5,500,077 Method of polishing/flattening diamond 5,499,601 Method for vapor-phase synthesis of diamond 5,496,596 Methods for producing diamond materials with enhanced heat conductivity 5,495,126 Polycrystalline diamond heat sink having major surfaces electrically insulated from each other 5,492,770 Method and apparatus for vapor deposition of diamond film
572
Malshe and Brown
5,491,002 Multilayer CVD diamond films 5,490,002 Multilayer CVD diamond films 5,490,963 Preparation of thin freestanding diamond films 5,488,350 Diamond film structures and methods related to same 5,488,232 Oriented diamond film structures on nondiamond substrates 5,487,945 Diamond films on nondiamond substrates 5,485,804 Enhanced chemical vapor deposition of diamond and related materials 5,480,686 Process and apparatus for chemical vapor deposition of diamond films using water-based plasma discharges 5,479,875 Formation of highly oriented diamond film 5,479,874 CVD diamond production using preheating
ACKNOWLEDGMENTS The authors would like to acknowledge Defense Advanced Research Projects Agency (DARPA) for their financial support. Discussions with Dr. Gordon, Dr. Naseem, and Dr. Schaper at High Density Electronics Center (HiDEC), University of Arkansas, Fayetteville, have been very helpful at many stages of research and are sincerely appreciated. Finally, the authors appreciate the efforts of Dr. Railkar of Materials and Manufacturing Research laboratory (MRL) at the University of Arkansas, Fayetteville in compiling this chapter.
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JA Carey, D Collins, BP Loh. Abstract from the IMAPS Advanced Technology Workshop on Thermal Spreaders and Heat Sinks, Beaver Creek, CO, April 1997. M Nagai, K Tanabe, Y Kumazawa, N Fujimori. Presented at the ISHM-97 Conference, Philadelphia, 1997. J Crumpton, R Koba, K Valenta, B Speck, C Roach, D Kuty, R Keusseyan. Abstract submitted to the IMAPS Advanced Technology Workshop on Thermal Spreaders and Heat Sinks, Beaver Creek, CO, April 1997. D Tilton. Electron Packaging Prod August:71, 1995. A Feldman, HPR Frederikse. Proceedings of Second International Conference on the Applications of Diamond Films and Related Materials, Omiya Saitama, Japan, August 25–27, 1993, p 261. AP Malshe, BS Park, WD Brown, A Muyshondt, HA Naseem, Diamond Relat Mater 8:1198, 1999.
12 Diamond Active Electronic Devices Hiromu Shiomi Sumitomo Electric Industries Ltd., Itami, Japan
I.
INTRODUCTION
Diamond is a promising material for electronic and optical devices under severe conditions because it has a large electron energy gap. Diamond is also an attractive material for high-frequency and high-power devices due to its high breakdown field, high electron and hole mobilities, and low dielectric constant. Figures of merit for materials are useful to compare diamond and other semiconductors and to consider the advantages of different materials for certain types of electronic applications. The author shows three figures of merit, Johnson’s figure of merit, Keyes’ figure of merit, and Baliga’s figure of merit, and discusses the performance of diamond devices. Johnson’s figure of merit [1] is usually used to compare material properties for power microwave applications. It is given by JFM =
冉 冊 EMvs 2
2
(1)
where vs is the scattering limited saturated velocity of carriers in the semiconducting material and EM is the peak electric field strength at breakdown. Diamond’s high electron saturation velocity (vs = 2.8 ⫻ 107 cm/sec at room temperature) [2] combined with a breakdown field EM in excess of 107 V/ cm [3] results in a high Johnson figure of merit, JFMdiamond/JFMSi = 2601. Keyes’ figure [4] of merit is an indication of the thermal limitation on a material’s high-frequency electrical performance. It is given by KFM =
冉 冊 cvs 4r
1/2
(2) 579
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Shiomi
where c is the velocity of light, r is the dielectric constant, and is the thermal conductivity. Diamond’s high electron saturation velocity and high thermal conductivity make Keyes’ figure of merit high, for example, KFMdiamond/KFMSi = 32. Baliga et al. [5,6] proposed a practical figure of merit for power devices. Although higher values of JFM and KFM are desirable for a good power semiconductor switch, they are not sufficient criteria. From a practical point of view, the drift region resistance is an important parameter that needs to be minimized for increasing the power-handling capability of a semiconductor device. Baliga’s figure of merit is based on the parameter A, which can be interpreted as the on-state drift region conductance per unit area of an abrupt one-sided p-n junction, A was given by Shenai et al. [5],
A =
qNB WM
(3)
where is the low-field carrier mobility, NB is the drift region doping density, WM is the depletion layer depth at breakdown, and q is the electronic charge. This equation was later corrected for semiconductors with deep donors or acceptors by Collins [7],
A =
qnc WM
(4)
where nc is the concentration of charge carriers. Assuming an abrupt one-sided p-n junction on uniformly doped material, the drift region doping density can be related to avalanche breakdown as NB =
sE2M 2qVB
(5)
where s is the permitivity of the semiconductor and EM is the peak electric field strength at breakdown. The depletion layer depth at breakdown, WM, is 2VB/EM. For the calculation of A in Eq. (4), nc = NB for Si and GaAs at or above room temperature. However, this is not the case for diamond. The ionization energy of boron acceptors in diamond is 0.37 eV [8], and only about 0.2% of the acceptors are ionized at room temperature in a typical natural semiconducting diamond [9]. The following expression must be used to determine the hole concentration p: p(p ⫹ ND) = (NA ⫺ ND ⫺ p)
冉
冊
2m*kT h2
3/2
exp
冉 冊 ⫺EA kT
(6)
where NA is the boron concentration, ND is the donor concentration, and m* is the hole density-of-states effective mass [9].
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581
The author calculated the temperature dependence of A as already described for a one-sided junction made of diamond, Si, GaAs, and 6H-SiC, as shown in Fig. 1, using the parameters after Collins [7]. Collins used an activation energy of 0.37 eV and a low value for the mobility of p-type diamond, 290 cm2 V⫺1 sec⫺1, which was the typical value for homoepitaxial diamond films. For these parameters, A is larger for diamond than for GaAs or Si at temperatures greater than about 180⬚C because of the deep level of the acceptor. However, the actual carrier concentrations are expected to be more because the energy level EA decreases as the doping concentration increases [10–12]. In the temperature range investigated [10], the activation energy changed from 0.2 to 0.37 eV due to the different acceptor concentrations. After modification of the growth process, mobilities between 1000 and 1430 cm2 V⫺1 sec⫺1 were achieved in the improved chemical vapor deposition (CVD) films, which had B concentrations between ⬃5 ⫻ 1016 and 2.2 ⫻ 1018 cm⫺3 [13]. If the mobility is 1300 cm2 V⫺1 sec⫺1 with an activation energy of 0.37 eV, A is the larger for diamond than for GaAs or
Figure 1 Conductance/cm2 calculated for one-sided diodes made from diamond, SiC, GaAs, and Si.
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Shiomi
Si at temperatures greater than about 110⬚C. When the author uses an activation energy of 0.22 eV [10,14] and a mobility of 1300 cm2 V⫺1 sec⫺1, A is larger for diamond than for GaAs or Si in the entire temperature range and for 6H-SiC at temperatures greater than about 70⬚C. It is concluded that diamond-based devices offer significant advantages in high-power applications. In the case of a high-frequency system, Trew et al. [15] examined the frequency performance of GaAs, silicon carbide, and diamond electronic devices, as shown in Fig. 2. Because frequency performance scales directly with saturation velocity, diamond devices are expected to operate at higher frequencies than GaAs or silicon carbide devices. Diamond MESFETs are capable of producing over 200 W of X-band power as compared with about 8 W for GaAs and about 60 W for silicon carbide. Trew’s group [16] further analyzed the output performance of diamond MESFETs and showed its
Figure 2 RF power performance versus frequency for diamond, SiC, and GaAs MESFETs. (From Ref. 15.)
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583
temperature dependence. Power performance of the diamond MESFETs exceeds that of GaAs MESFETs when operated at temperatures higher than 177⬚C. The author speculates that the temperature limit should be lower, following a discussion similar to the earlier discussion of on-state drift region conductance.
II.
DIAMOND TRANSISTORS
A.
Introduction
Table 1 shows the pioneering works on diamond transistors. The author categorizes the history of the development of diamond transistors into three periods: 1982–1988 for preliminary bulk transistors, 1989–1992 for preliminary field-effect transistors, and 1993– for improved FETs and circuits.
Table 1
Historical Development of Diamond Transistors
Year
Topic
1982 1987
Bipolar transistor on natural diamond Point-contact transistor on HPHT diamond Permeable base transistor on HPHT diamond First MESFET on CVD diamond films First MESFET on CVD diamond films MiSFET with CVD intrinsic diamond MOSFET with SiO2 MESFET by rapid thermal processing MOSFET by the selective growth method MOSFET with ion implanted channel MiSFET on polycrystalline diamond films Gate-recessed MOSFET MOSFET on polycrystalline diamond films First diamond circuit MESFET using surface conductive layer Self aligned gate-recessed MESFET MOSFET on improved diamond films Pulse-doped (␦-doped) MESFET Normally-off diamond MESFET Transconductance exceeding 10 mS/mm
1988 1989 1990
1991
1992
1993 1994
1995 1997 1999
Researchers
Ref.
F. Prins W. Geis et al.
17 19
W. Geis et al.
20
Shiomi et al. Sh. Gildenblat et al. Shiomi et al. A. Grot et al. Tsai et al. Sh. Gildenblat et al.
22 25 26 27 30 28
R. Zeisse et al. Nishimura et al.
32 34
A. Grot et al. J. Tessmer et al.
29 37
Zeidler et al. Kawarada et al. Shiomi et al. L. Dreifus et al. Shiomi et al. J. Looi et al. Tsugawawa et al.
39 41 44 45 47 48 49
584
B.
Shiomi
1982–1988, Preliminary Bulk Transistors
The first active diamond electronic device was demonstrated in 1981 when Prins [17] reported bipolar transistor action in ion-implanted natural diamond. The device was fabricated by ion bombarding C⫹ ions onto the surface of a natural p-type diamond in two regions separated by 3.2 m, as shown in Fig. 3a. The bombarded regions served as the n-type regions, although they were actually radiation-damaged regions. The common emitter current gain of the device is less than one, although the output I-V characteristics do show complete current saturation, as shown in Fig. 3b. In 1987, Tzeng et al. [18] reported a similar npn bipolar junction transistor, fabricated by ion implanting arsenic in a p-type natural diamond. The current gain obtained in the device was about 0.8. A point-contact transistor, operational at 510⬚C, was reported by Geis et al. [19]. This is the first report of diamond transistors that have power gain. Point-contact transistors were fabricated on a synthetic boron-doped diamond produced by a high-pressure, high-temperature (HPHT) process. From room temperature capacitance-voltage measurements of Schottky diodes on the diamond, it was determined that the net ionized boron acceptor concentration was on the order of 1016 cm⫺3. The base, collector, and emitter contacts consisted of tungsten probes that formed Schottky or ohmic contacts on the diamond surfaces, as shown in Fig. 4a. The collector current as a function of base-collector voltage is shown in Fig. 4b. When operated at room temperature, the small-signal current gain (the ratio of the collector to emitter currents) varied from 2 to 25 and the calculated small-signal power gain varied from 6 to 35, depending on the transistor and the operating conditions. At a high temperature, 510⬚C, the current gain typically varied from 1.6 to 0.5 and the power gain varied from 4.5 to 1.3, depending on the devices. In the following year, Geis et al. [20] reported the fabrication of a permeable base transistor. An ion-beam etching technique was used to fabricate the vertical structure of this device [21]. A schematic diagram of the device structure and I-V characteristics are shown in Fig. 5. A transconductance of 30 S/mm was obtained, which was limited by a parasitic resistance (⬃104 ⍀ cm) in the substrate of the device, not an inherent material property nor the structure.
C.
1989–1992, Preliminary Field-Effect Transistors
In 1989, the first diamond MESFETs were fabricated independently by two groups on boron-doped homoepitaxial diamond films grown by the plasmaassisted chemical vapor deposition method [22–25]. Shiomi et al. [22] ep-
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585
Figure 3 First active diamond electronic device, an npn bipolar junction transistor in natural p-type diamond. (a) Schematic illustration of the masking of the base area during carbon ion implantation. (b) I-V characteristics. (From Ref. 17.)
586
Shiomi
Figure 4 Diamond point-contact transistor. (a) Schematic drawing of the transistor. The ohmic contact was formed either by tungsten probes or by e-beam evaporated tungsten on the (111) facets of the diamond. (b) I-V characteristics. (From Ref. 19.)
itaxially grew diamond films on an HPHT diamond (100) surface in the selected condition of high CH4 concentration with B2H6 as dopant gas. Ti was used for the source and drain ohmic contacts and Al was used for the gate Schottky contacts. The gate length and width were large, 140 m and 1.8 mm, respectively, because the MESFET patterns were fabricated directly with a shadow mask without using a photolithography technique. The device structure and the drain current versus drain voltage characteristics are shown in Fig. 6. The gate voltage was varied from ⫺1.0 to 5.0 V when the leakage current was less than 1 A. The gate voltage controlled the drain current
Diamond Active Electronic Devices
587
Figure 5 Diamond permeable base transistor. (a) Device structure. To fabricate this device, ion beam-assisted etching was used to produce the grating in the diamond. (b) I-V characteristics. (From Ref. 20.)
588
Shiomi
Figure 6 First diamond MESFET. (a) Device structure. (b) I-V characteristics. (From Ref. 22.)
although the ohmic property of the source to drain was not optimized without annealing due to the fabrication process. Gildenblat et al. [25] fabricated a MESFET with a circular Al gate, as shown in the inset of Fig. 7. The leakage current remained under 50 nA at reverse biases up to 16 V. This structure shows some degree of gate control over the drain current ID. As can be seen in Fig. 7, a 6 V reverse gate bias results in a reduction of the drain current. Following the first fabrication of MESFETs, the two groups previously mentioned made attempts to improve the device performance by reducing the gate leakage with the incorporation of a gate insulator. A field-effect transistor with an insulated gate is generally called an insulated-gate field-
Diamond Active Electronic Devices
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Figure 7 First diamond MESFET; device layout in the inset and I-V characteristics. (From Ref. 25.)
effect transistor (IGFET) or a metal-insulator-semiconductor field-effect transistor (MISFET). Shiomi et al. [26] inserted a 0.23-m-thick intrinsic diamond layer between the gate metal and boron-doped diamond layer. To specify this structure in Table 1, the author calls it a metal–intrinsic semiconductor–semiconductor (MiS) structure. The device structure and drain current versus drain voltage characteristics are shown in Fig. 8. The gate voltage was varied from 0 to 20 V in 5 V steps and the drain voltage was applied up to ⫺50 V. The drain current tends to saturate as the drain voltage increases. The intrinsic diamond layer withstood 70 V between gate and drain. The maximum transconductance was 2.0 S/mm. Grot et al. [27,28] used 0.10-m-thick SiO2 as a gate insulator. It was formed by sputtering in an Ar (50%)–O2 (50%) plasma at 350⬚C. To specify this device structure in Table 1, the author calls it a metal-oxide-semiconductor (MOS) structure. During the fabrication of these devices, they used a selective growth process, where boron-doped homoepitaxial diamond was grown in the opening of sputtered SiO2. The hole mobility of selectively grown films varied between 210 and 290 cm2/V sec for hole concentrations between 1.0 ⫻ 1014 and 6.9 ⫻ 1014 cm⫺3. The device structure is shown in Fig. 9a. An Au (1.25 m)/Ti (50 nm) metallization was deposited on top of the selectively grown homoepitaxial diamond island. The transistor characteristics at 300⬚C are shown in Fig. 9b. High-temperature operation of the device is made possible by incorporation of an insulating gate. The gate leakage current remained under 10 pA even during operation at 300⬚C. To
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Figure 8 First diamond MiSFET with intrinsic diamond layer between gate metal and boron-doped diamond layer. (a) Device structure. (b) I-V characteristics, illustrating the tendency of drain current saturation at high positive gate voltage bias. (From Ref. 26.)
improve the transistor characteristics, Grot et al. [29] introduced the electron cyclotron resonance (ECR) plasma etching process and fabricated mesaisolated recessed gate field-effect transistors that operated at temperatures up to 350⬚C. The maximum transconductance was 87 S/mm at 200⬚C, although they did not exhibit complete saturation or pinch-off of the channel. In 1990, Tsai et al. [30,31] reported on the fabrication of a diamond MESFET with the channel formed by solid-state diffusion. The device structure and the I-V characteristics are shown in Fig. 10. Boron was diffused into type IIa diamonds and simultaneously electrically activated by a rapid thermal annealing technique using a cubic boron nitride planar diffusion source in an argon atmosphere. The annealing took place at ⬃1370⬚C for 20 sec. Boron concentrations as high as 3.5 ⫻ 1019 cm⫺3 were detected at ˚ in the diamond substrate using secondary ion mass speca depth of 500 A
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Figure 9 First diamond MOSFET with SiO2 as a gate insulator. (a) Device structure. (b) I-V characteristics at 300⬚C. (From Ref. 27.)
trometry (SIMS). Electrical contacts of Ti/Au were made on diamond via evaporation, and subsequent ohmic annealing was carried out for 30 min at 800⬚C. Finally another, unannealed Ti/Au bilayer 0.15 mm in length and 1 mm in width was deposited to form the Schottky gate. The MESFET was observed to pinch off at high positive gate bias, and reverse-bias leakage was below 10⫺12 A at 5 V. In addition to the boron doping processes already mentioned—the epitaxial growth process and the solid-state diffusion process—Zeisse et al. [32] introduced the ion-implantation process and fabricated MOSFETs. The device structure and the I-V characteristics are shown in Fig. 11. A p-type conducting layer was formed in a substrate of semi-insulating natural dia-
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Figure 10 Diamond MESFET with ultrashallow boron-doped layer formed by rapid-thermal-processing (RTP) solid-state diffusion using cubic boron nitride as the dopant source. (a) Device structure with an ultrashallow p-doped channel less than ˚ at a boron concentration of about 3 ⫻ 1019 cm⫺3. (b) I-V characteristics, 500 A illustrating pinch-off behavior at high positive gate bias. (From Ref. 31.)
mond (type IIa) by boron implantation. Boron ions were implanted at 80 K using a multiple implant scheme (25 keV, 1.5 ⫻ 1014 B⫹/cm2; 50 keV, 2.1 ⫻ 1014 B⫹/cm2; and 100 keV, 3.0 ⫻ 1014 B⫹/cm2) intended to provide a uniform p-type layer 210 nm thick. After implantation, the diamond was annealed at 1263 K in nitrogen to remove the implantation damage and
> Figure 11 Diamond MOSFET fabricated by ion implanting boron into natural insulating type IIa diamond. (a) Plan (top) and cross section (bottom) views of an ion-implanted MOSFET. The thickness of the implanted layer is approximately 210 nm. The circular geometry does not require a mesa etch to contain the current flow from source to drain. (b) I-V characteristics at room temperature. The transconductance approaches 9.5 S toward saturation for a change in gate voltage from 0 V to 2 V. (From Ref. 32.)
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activate the implanted boron. The gate insulator, consisting of an SiO2 film approximately 100 nm thick, was deposited by indirect plasma-enhanced chemical vapor deposition [33] at a temperature of 300⬚C. Saturation and pinch-off were both observed at room temperature. The transconductance was 3.9 S/mm and the output conductance was 60 nS/mm. This was the first reported use of ion implantation to fabricate a field-effect device successfully in diamond. The fabrication of active devices on polycrystalline diamond films is attractive because of the availability of large-area inexpensive substrate material. Nishimura et al. [34–36] first reported field-effect transistor action in polycrystalline diamond films. The device structure and the I-V characteristics are shown in Fig. 12. The maximum transconductance was 5 S/mm. The device did not show saturation or pinch-off. Tessmer et al. [37,38] improved the polycrystalline diamond FETs and obtained a larger estimated maximum transconductance of 121 S/mm, although the device did not show the pinch-off condition.
D.
1993– , Improved FETs and Circuits
Zeidler et al. [39] fabricated a simple demonstration circuit with a voltage gain of two using two MOSFETs. Natural type IIa diamonds were implanted with boron at 77 K, using the multiple scheme, low-temperature boron ion implantation procedure to produce an approximately uniformly doped p-type layer about 200 nm thick. This layer was used to fabricate MOSFETs operating in the depletion mode, with good uniformity from device to device. Two of the devices made the first diamond electric circuit, as shown in Fig. 13a. As indicated by the labels ‘‘DOT’’ and ‘‘RING’’ in Fig. 13a, the transistors were connected in such a way as to place a parasitic resistor across the power supply rather than across the output. The circuit was tested at room temperature under direct-current conditions. A gain (Vout/Vin) of approximately two was measured over a wide range of VDD, as seen in Fig. 13b. Kawarada et al. [40,41] fabricated MESFETs using a surface conductive layer of hydrogen-terminated undoped homoepitaxial diamond films. Hydrogen-terminated films display p-type conduction [10,42,43]. These FETs exhibited enhancement-mode effects and a high value of transconductance of 200 S/mm (Fig. 14), although these operations should depend on the environment, which controls the surface conductive layer. Shiomi et al. [44] introduced a reactive ion etching process and fabricated MESFETs with a self-aligned recessed structure. The device structure and the I-V characteristics are shown in Fig. 15. Saturation and pinch-off characteristics were both observed in diamond MESFETs for the first time.
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Figure 12 Polycrystalline diamond transistor. (a) Top and cross-sectional views of the MiSFET. (b) I-V characteristics of the MiSFET at room temperature. The inset shows the I-V characteristics between source and gate. (From Ref. 36.)
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Figure 13 Diamond circuit. (a) Circuit configuration with parasitic resistances shown in dashed lines. (b) Voltage gain as a function of applied bias. (From Ref. 39.)
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Figure 14 Diamond MESFET using a surface conductive layer of undoped homoepitaxial film. (a) Cross section of the MESFET. (b) I-V characteristics of the MESFET with a gate length of 10 m. The transconductance obtained is 200 S/mm. (From Ref. 41.)
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Figure 15 Gate-recessed structure MESFET. (a) Cross section of MESFET. (b) IV characteristics of the MESFET with a gate length of 4 m. (From Ref. 44.)
Dreifus et al. [13,45,46] demonstrated diamond MOSFETs that exhibited complete pinch-off of the active channel and a large transconductance on the order of a mS/mm at the elevated temperature. In this research, highquality homoepitaxial films of semiconducting diamond were deposited on IIa diamond substrates. The room temperature hole mobility exceeded 1400 cm2/V sec. In addition, compensation was measured as 3 ⫻ 1015 cm⫺3, which was similar to donor levels in the best natural type IIb single crystal diamonds. These films were used to fabricate MOSFETs. Pinch-off and saturation of the channel current were observed at temperatures up to 500⬚C. Figure 16 shows the I-V characteristics at 400⬚C. Diamond field-effect transistors were configured in analog-amplifier and a two-input NOR digital logic circuit. Successful logic operation has been observed up to 400⬚C. Shiomi et al. [47] have fabricated MESFETs utilizing a boron pulsedoped layer as a conducting channel. The layered structure of the MESFET, as shown in the inset of Fig. 17a, consists of an undoped 150-nm-thick buffer layer grown first, followed by a boron pulse-doped layer of 11 nm (the full
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Figure 16 Diamond MOSFET with an improved epitaxial layer; I-V characteristics at 400⬚C. Gate-to-source voltage varies from 0 to 14 V in 2 V increments. (From Ref. 13.)
width of the half-maximum of the boron profile in Fig. 17a) and a 75-nmthick undoped diamond cap layer. These layers were grown by controlling the flow of the dopant gas during the microwave plasma-assisted chemical vapor deposition. The doping levels in the undoped layers and the boron pulse-doped layer were found to be 1019 cm⫺3, respectively, as shown in Fig. 7a. The I-V characteristics in Fig. 7b showed that this fieldeffect transistor with a gate length of 4 m and a gate width of 39 m exhibited an extrinsic transconductance of 116 S/mm with both pinch-off characteristics and current saturation. In 1997, Looi et al. reported the first normally off enhancement mode MESFET structures formed with a near-surface p-type hydrogenated diamond layer [48]. Such devices play an enabling role in circuit design. The transistors were formed on free-standing polycrystalline diamond grown by microwave plasma-enhanced CVD. Thermally evaporated gold was used for the ohmic contacts and aluminum was used for the gate metal. With no gate bias, the drain-to-source current was less than 1 nA, corresponding to an off-state. By applying negative gate voltage, an on-state was established. With a gate voltage of ⫺1.8 V, the drain-to-source current magnitude was 4.7 A, with the gate current remaining less than 10 nA. Room temperature transconductance values were 0.14 mS/mm. On single crystalline diamond, a hydrogenated p-type layer can be utilized to form FET’s with even higher performance figures of merit. Tsugawa et al. formed both MESFET and MOSFET devices on homoepitaxial CVD diamond deposited on a synthetic Ib diamond (100) surface [49]. By
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Figure 17 Pulse-doped diamond p-channel MESFET. (a) Depth profile of boron concentration measured by SIMS and cross-sectional drawing of a pulse-doped MESFET in the inset. (b) I-V characteristics. (From Ref. 47.)
appropriate selection of gate metal, both normally-off (enhancement mode) and normally-on (depletion mode) devices were achieved. In this case, lead was used as the gate metal for an enhancement mode MESFET and copper was used as the gate metal for a depletion mode MESFET. Transconductance values exceeded 10 mS/mm for both enhancement mode and depletion mode transistors.
Diamond Active Electronic Devices
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601
SUMMARY AND DIRECTION OF FUTURE WORK
A simple analysis based on the measured peak electric field strength at avalanche breakdown is presented to evaluate the performance of various semiconducting materials for high-power electronics. The calculated results indicate that diamond is superior material for fabricating power semiconductor devices. The high saturation velocity of diamond is shown to enable diamond to handle high power at high frequency. Moreover, the highest thermal conductivity and large energy gap of diamond offer an insulating layer which can be conveniently used for conducting the heat away from the active region of devices. The development of transistors is a benchmark for a new semiconductor material. It has been made clear by many researchers in the last several years that diamond is not just a simple gemstone but also a semiconducting material. However, researchers have now also demonstrated the feasibility of diamond active devices. For practical application, the author thinks that we need major breakthroughs in the following three techniques; the growth technique, the doping technique, and the microfabrication technique. In terms of the growth technique, the high-speed growth technique or the heteroepitaxial growth technique should be developed in the quest to obtain large single-crystal diamond wafers or films. For example, high-quality 6H-SiC, which is supplied as a wafer by the improved sublimation method [50] and is grown as a film by step-control epitaxy [51], has accelerated the development of SiC high-power devices [52–55]. High-quality GaN has been heteroepitaxially grown on sapphire [56,57], and soon after the brightest blue light–emitting diode (LED) was successfully fabricated [58]. These examples show that the production of high-quality wafers or films is the main breakthrough for newcomer semiconductors. In terms of the doping technique, the concentration and distribution of dopant should be well controlled. And n-type diamond with low resistivity should be obtained for a wide range of electrical applications. The recent progress in improving CVD diamond may enable the growth of n-type diamond. In terms of the microfabrication technique, the processing tools for diamond devices must be further developed to the point of fabricating reproducible devices. Recent progress shows that microfabrication techniques established for other semiconductors can be utilized for diamond. At present, it is not clear that diamond active devices have performance competitive with that of conventional materials. However, not a year passes without significant progress in the material and device quality of diamond transistors. In the next few years, high-power and high-frequency diamond devices may surpass the performance of devices fabricated in conventional
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semiconductors and SiC. And the mentioned breakthroughs will open the way to the commercial use of diamond active devices. REFERENCES 1. 2. 3.
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EO Johnson. Physical limitations on frequency and power parameters of transistors. RCA Rev 26:163, 1963. DK Ferry. High field transport in wide-bandgap semiconductors. Phys Rev B 12:2361, 1975. AV Bogdanov, IM Vikulin, TV Bogdanova. Investigation of microplasma breakdown at a contact between a metal and a semiconducting diamond. Sov Phys Semicond 16:720, 1982. RW Keyes. Physical limits in digital electronics. Proc IEEE 63:740, 1975. K Shenai, RS Scott, BJ Baliga. Optimum semiconductor for high-power electronics. IEEE Trans Electron Devices ED-36:1811, 1989. BJ Baliga. Power semiconductor device figure of merit for high-frequency applications. IEEE Electron Device Lett 10:455, 1989. AT Collins. Diamond electronic devices—can they outperform silicon or GaAs? Mater Sci Eng B11:257, 1992. AT Collins, WS Williams. The nature of the acceptor centre in semiconducting diamond. J Phys C 4:1789, 1970. JE Field. The Properties of Diamonds. New York: Academic Press, 1979. H Shiomi, Y Nishibayashi, N Fujimore. Characterization of boron-doped diamond epitaxial films. Jpn J Appl Phys 30:1363, 1991. JW Glesener. Hole capture in boron-doped diamond. Appl Phys Let 64:217, 1994. M Werner, O Dorsch, HU Baerwind, E Obermeier, L Haase, W Seifert A. Ringhandt, C Johnston, S Romani, H Bishop, and PR Chalker. Charge transport in heavily B-doped polycrystalline diamond films. Appl Phys Lett 64:595, 1994. JT Glass, DL Dreifus, RE Fauber, BA Fox, ML Hartsell, RB Henard, JS Holmes, D Malta, LS Plano, AJ Tessmer, GJ Tessmer, HA Wynands. High quality semiconducting diamond films and devices. Fourth International Conference on New Diamond Science and Technology, MYU, Kobe, Japan, 1994, p 335. N Fujimori, H Nakahata, T Imai. Properties of boron-doped epitaxial diamond films. Jpn J Appl Phys 29:824, 1990. RJ Trew, J Yan, PM Mock. The potential of diamond and silicon carbide electronic devices for microwave and millimeter-wave power applications. Proc IEEE, 79:598, 1991. MW Shin, RJ Trew, GL Bilbro. High temperature dc and RF performance of p-type diamond MESFET: comparison with n-type GaAs MESFET. IEEE Electron Device Lett 15:292, 1994. JF Prins. Bipolar transistor action in ion implanted diamond. Appl Phys Lett 41:950, 1982.
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Y Tzeng, TH Lin, JL Davidson, LS Lan. Fabrication and high temperature characteristics of diamond electronic devices. Proceedings of Seventh Biennial University/Government/Industry Microelectronics Symposium, IEEE, Rochester, NY, 1987, p 187. MW Geis, DD Rathman, DJ Ehrlich, RA Murphy, WT Lindley. High-temperature point-contact transistors and Schottky diodes formed on synthetic borondoped diamond. IEEE Electron Device Lett EDL-8:341, 1987. MW Geis, NN Efremow, DD Rathman. Summary abstract: device applications of diamonds. J Vac Sci Technol A 6:1953, 1988. NN Efremow, MW Geis, DC Flanders, GA Lincoln, NP Economou. Ion-beam– assisted etching of diamond. J Vac Sci Technol B 3:416, 1985. H Shiomi, Y Nishibayashi, N Fujimore. Field-effect transistors using borondoped diamond epitaxial films. Jpn J Appl Phys 28:L2153, 1989. H Shiomi, Y Nishibayashi, N Fujimore. Diamond transistors. Fiftieth Fall Meeting of The Japan Society of Applied Physics and Related Society, Japan, 28a-E-10, 1989, p 384. N Fujimori, T Imai, H Nakahata, H Shiomi, Y Nishibayashi. Epitaxial growth of diamond and diamond devices. Proceedings of Fall 1989 MRS Meeting, Boston, Vol 162, 1989, p 23. GS Gildenblat, SA Grot, CW Hatfield, CR Wronski, AR Badzian, T Badzian, R Messier. Electrical properties of homoepitaxial diamond films. Proceedings of Fall 1989 MRS Meeting, Boston, Vol 162, 1989, p 297. H Shiomi, Y Nishibayashi, N Fujimori. Characterization of boron-doped diamond epitaxial films and applications for high-voltage Schottky diodes and MESFETs. Second International Conference of New Diamond Science and Technology, Material Research Society, Washington, DC, 1990, p 975. SA Grot, CW Hatfield, GS Gildenblat, AR Badzian, T Badzian. Semiconductor device development using selectively grown thin-film diamond. Second International Conference of New Diamond Science and Technology, Washington, DC, 1990, p 949. GS Gildenblat, SA Grot, AR Badzian. High-temperature thin-film diamond field-effect transistor fabricated using a selective growth method. IEEE Electron Device Lett EDL-12:37, 1991. SA Grot, GS Gildenblat, AR Badzian. Diamond thin-film recessed gate fieldeffect transistors fabricated by electron cyclotron resonance plasma etching. IEEE Electron Device Lett EDL-13:462, 1992. W Tsai, M Delfino, LY Ching, G Reynolds, D Hodul, CB Cooper III. Boron doping of diamond via solid state diffusion. Second International Conference of New Diamond Science and Technology, Material Research Society, Washington, DC, 1990, p 937. W Tsai, M Delfino, D Hodul, M Riaziat, LY Ching, G Reynolds, CB Cooper III. Diamond MESFET using ultrashallow RTP boron doping. IEEE Electron Device Lett EDL-12:157, 1991. CR Zeisse, CA Hewett, R Nguyen, JR Zeidler, RG Wilson. An ion-implanted diamond metal-insulator-semiconductor field-effect transistor. IEEE Electron Device Lett EDL-12:602, 1991.
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Shiomi LG Meiners. Indirect plasma deposition of silicon dioxide. J Vac Sci Technol 21:655, 1982. K Nishimura, R Kato, S Miyauchi, K Kobashi. Metal semiconductor field effect transistor with polycrystalline diamond films. Japan New Diamond Forum, Fifth Diamond Symposium, Tsukuba, Japan, 1991, p 34. K Nishimura, H Koyama K, Miyata, K Suzuki, K Kobashi. Fabrication of field effect transistors made of polycrystalline diamond films. Third International Symposium on Diamond Materials at the ECS Spring Meeting, Hawaii, 1993. K Nishimura, K Kumagai, R Nakamura, K Kobashi. Metal/intrinsic semiconductor/semiconductor field effect transistor fabricated from polycrystalline diamond films. J Appl Phys 76:8142, 1994. AJ Tessmer, K Das, DL Dreifus. Polycrystalline diamond field-effect transistors. Diamond Relat Mater 1:89, 1992. AJ Tessmer, LS Plano, DL Dreifus. High-temperature operation of polycrystalline diamond field-effect transistors. IEEE Electron Device Lett EDL-14:66, 1993. JR Zeidler, CA Hewett, R Nguyen, CR Zeisse, RG Wilson. A diamond driveractive load pair fabricated by ion implantation. Diamond Relat Mater 2:1341, 1993. N Nakamura, M Aoki, M Itoh, N Jin, H Kawarada. Fabrication of MESFET utilizing hydrogen-terminated homoepitaxial diamond films. Fourth International Conference on New Diamond Science and Technology, MYU, Kobe, Japan, 1994, p 729. H Kawarada, M Aoki, M Ito. Enhancement mode metal-semiconductor field effect transistors using homoepitaxial diamonds. Appl Phys Lett 65:1563, 1994. MI Landstrass, KV Ravi. Resistivity of chemical vapor deposited diamond films. Appl Phys Lett 55:975, 1989. H Nakahata, T Imai, N Fujimore. Change of resistance of diamond surface by reaction with hydrogen and oxygen. Second International Symposium on Diamond Materials, The Electrochemical Society, Washington, DC, Vol 91, 1991, p 487. H Shiomi, Y Nishibayashi, N Toda, S Shikata, N Fujimori. Fabrication of a gate-recessed-structure MESFET on diamond films. Fourth International Conference on New Diamond Science and Technology, MYU, Kobe, Japan, 1994, p 661. DL Dreifus, AJ Tessmer, JS Holmes, C-T Kao, DM Malta, LS Plano, BR Stoner. Diamond field-effect transistors. MRS Spring Meeting, San Francisco, 1994. BA Fox, ML Hartsell, DM Malta, HA Wynands, C Kao, LS Plano, GJ Tessmer, RB Henard, JS Holmes, AJ Tessmer, DL Dreifus. Diamond devices and electrical properties. Diamond Relat Mater 4:622, 1995. H Shiomi, Y Nishibayashi, N Toda, S Shikata. Pulse-doped diamond p-channel metal semiconductor field-effect transistor. IEEE Electron Device Lett EDL16:36, 1995.
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13 Diamond Film Optics D. K. Reinhard Michigan State University, East Lansing, Michigan
I.
INTRODUCTION
The optical properties are among the most interesting of the physical properties of diamond. These include the widest spectral optical transmission range of all known solid materials [1] and, for a transparent material, an unusually high index of refraction. Such attributes, combined with extreme hardness, high thermal conductivity and chemical resistance, make diamond a material of obvious interest for optical applications. However, optical applications of diamond have traditionally been very limited because of the notoriously high cost of quality diamond, even for samples of modest dimensions. In spite of this, early experiments brought convincing evidence of the potential of diamond for optical uses. A case in point is the diamond window used as the optical port for an infrared radiometer experiment on the Pioneer Venus probe launched in 1978. This 18.2 mm diameter by 2.8 mm thick diamond window survived earth atmosphere on launch, the cold and vacuum of space, and the Venus atmosphere (largely carbon dioxide with various acids, including sulfuric acid, at temperatures of 800 K and pressures of 90 atm) [1]. Other important demonstrations included the use of single-crystal diamond as high-power laser windows. By virtue of low optical absorption combined with high thermal conductivity, the windows successfully passed CO2 laser beams with continuous wave (CW) power densities in excess of 1 MW/cm2 [2]. Such experiments demonstrated unique optical capabilities for diamond; remaining as a significant barrier for more widespread use was the prohibitive economic cost. 607
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The development of chemical vapor deposition (CVD) methods to synthesize diamond dramatically altered the issue of material availability. CVD technology has demonstrated the ability to create diamond structures of sizes that would be prohibitively expensive or simply not available in the traditional single-crystal form. For example, CVD has been used to form freestanding diamond structures in a variety of shapes including domes, tubes, and plates [3] as well as to form thin-film diamond coatings on a variety of substrates with diameters up to 20 cm and larger. Consequently, optical applications not previously feasible are now within consideration. Such applications include freestanding optical elements as well as coatings of lenses and other optical elements and devices. Nevertheless, there are several nontrivial issues associated with the application of CVD diamond for optical purposes. The intrinsic optical properties of diamond are based upon high-purity, single-crystalline samples. CVD samples, however, are generally polycrystalline and, depending on the deposition processes, may or may not contain significant defects and impurities. An additional concern with CVD material is mechanical strength and, of particular issue for thin-film applications, the integrity of a diamond coating in terms of film stress and adhesion to the substrate. This chapter is concerned primarily with the optical properties associated with polycrystalline CVD diamond and with related optical application issues of both thick films and thin films. The optical properties of singlecrystal diamond are first reviewed briefly in order to provide a frame of reference. Next, optical properties of freestanding polycrystalline films are considered, followed by a discussion of diamond film–coated substrates. Finally, electro-optic applications of diamond are briefly considered, including photodetection, photoconductive switching, and photon emission. The scope of the chapter is primarily concerned with the ultraviolet, visible, and infrared, with only passing reference to other portions of the electromagnetic spectrum. Regarding the illustrations and data presentation in this chapter, it is noted that the spectrally dependent optical properties of materials are variously presented as a function of wavelength, , or as a function of wavenumber (that is, ⫺1, or inverse wavelength), or as a function of photon energy, E, which is related to the free-space wavelength as E = hc/, where h is Planck’s constant and c is the velocity of light in free space.* All three methods of data presentation are in common use in the literature and are used in this chapter.
*If the energy is expressed in electron volts and the free-space wavelength in micrometers, then E = 1.24/.
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OPTICAL PROPERTIES OF SINGLE-CRYSTAL DIAMOND
A.
Reflection and Absorption
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Optical phenomena associated with single-crystal diamond have been a topic of study for hundreds of years. Such studies have been motivated in significant part by the fact that diamond gem commerce is based on the optical properties of naturally occurring single-crystal diamond as obtained from the earth. These properties have been extensively reviewed, for example, by Davies [4] and by Clark [5]. They are discussed here briefly as a prerequisite to the discussion of polycrystalline CVD diamond. Consider an optical beam of intensity I0 in air normally incident on the surface of a diamond sample that is characterized by an optical absorption coefficient ␣. The optical intensity at some depth d in the sample, neglecting reflections from the back of the sample, is I(d) = I0(1 ⫺ R)exp(⫺␣d)
(1)
where R is the reflection coefficient at the front surface given by R=
冉
na ⫺ nd na ⫹ nd
冊
2
(2)
na and nd being the refractive indices of air and diamond, respectively. The refractive index of diamond is given by the Sellmeier equation [6], n2d ⫺ 1 =
4.33562 0.33062 2 ⫹ 2 ⫺ (0.1060) ⫺ (0.1750)2 2
(3)
where is the wavelength expressed in m. Figure 1 illustrates the wavelength dependence of the refractive index. There is clearly considerable dispersion in the visible portion of the spectrum, with nd ranging from near 2.47 in the violet to 2.40 in the red, a property that is also enhanced by the cut of diamond gems—light is refracted into a spectrum of colors when it passes out into the air and leads to the ‘‘fire’’ of diamond. Diamond, in fact, shows the largest color dispersion of the gems. The large refractive index of diamond causes a large normal incidence reflection value, approximately 17%, and also results in a small critical angle for total reflection, approximately 24 degrees.* Equation (3) is valid from the infrared to the onset of band-to-band absorption in the ultraviolet. The room-temperature diamond electronic band *This is taken advantage of by gem cutters. Facets in the so-called brilliant cut are arranged such that it is not possible for light to fall on the lower facets at lessor angles, so no light is refracted out the back.
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Figure 1 Wavelength dependence of the refractive index according to the Sellmeier equation.
gap, or separation between the valance-band maximum and conduction-band minimum [5], is 5.47 eV ( ⬇ 0.23 m), which effectively establishes an upper limit on photon energy for transmission.* Because the energy gap in diamond is indirect, photon excitation of an electron from the valence band to the conduction band must be accompanied by phonon absorption or emission in order to conserve momentum. For an indirect gap semiconductor with energy gap EG the absorption coefficient corresponding to transitions involving absorption of a photon of energy E and absorption of a phonon of energy EP is of the general form [7]
␣a =
A(E ⫺ EG ⫹ EP)2 exp(Ep/kT) ⫺ 1
(4)
for E > EG ⫺ EP and zero for E < EG ⫺ EP and where A is a slowly varying *According to Clark [5], the energy gap decreases at a rate of 5.4 ⫻ 10⫺5 eV/K.
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function of photon energy. For the emission of a phonon the absorption coefficient is
␣e =
A(E ⫺ EG ⫺ EP)2 1 ⫺ exp(⫺Ep /kT)
(5)
for E > EG ⫹ EP and zero for E < EG ⫹ EP. When the photon energy is sufficient to cause phonon emission, ␣e dominates ␣a. The actual situation is somewhat more complex in that more than one type of phonon must be considered* and excitons can also play a role. Clear from the general theory, however, is the fact that for photon energies greater than the energy gap, the absorption increases sharply. Experimental results for diamond show that the absorption coefficient increases from approximately 15 cm⫺1 at 5.5 eV to 5000 cm⫺1 at 5.96 eV [4,5]. In addition to the fundamental absorption described by Eqs. (4) and (5), many semiconductors are observed experimentally to exhibit an exponential increase in absorption with energy near the band gap. The absorption coefficient associated with this observation is modeled empirically by Urbach’s rule according to the expression [8]
␣U = ␣0 exp
冋
册
(T)(E ⫺ E0) kT
(6)
where T is the temperature in kelvins and k is Boltzmann’s constant. Experimental observation of such absorption has been reported for high-purity single-crystal diamond. Thomas and Tropf [9] reported Urbach tail parameters at room temperature to be ␣0 = 4.23 ⫻ 1011 cm⫺1 and = 0.585 with E0 being the value of the direct energy gap in diamond, 6.5 eV. Figure 2 illustrates the wavelength dependence of the Urbach edge absorption coefficient which contributes an exponential absorption edge tail below the band gap. In addition to the Urbach tail, a temperature-independent weak absorption tail is experimentally observed that extends farther into the band gap, from approximately 0.23 to 0.35 m. The magnitude of the weak tail varies somewhat from sample to sample [5] but on a representative high-purity single-crystal sample has been reported to be modeled as [9]
␣w = 0.28 exp(0.45E)
(7)
where E is in eV and the absorption coefficient is in cm⫺1. The weak tail absorption is primarily an ultraviolet phenomenon—the value of ␣w becomes on the order of 1 cm⫺1 by the blue portion of the visible spectrum. *Phonon energies at high-symmetry points in the Brillouin zone range in energy from 70 to 165 meV [4].
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Figure 2 Wavelength dependence of the absorption coefficient near the band edge as calculated by an Urbach edge model.
For longer wavelengths, throughout the visible and into the infrared, where the absorption coefficient as modeled in Eqs. (4)–(7) becomes zero or vanishingly small, high-purity diamond should exhibit essentially no optical absorption until excitation of vibration modes of C — C bonds occurs in the range of approximately 2 to 7 m. Diamond in the absence of significant defects or impurities has no infrared-active single-phonon absorption because diamond is a symmetric covalent material with no dipole moment to interact with the electric field of light. However, even in pure diamond, combinations of phonons can create weak dipole moments and two-phonon, three-phonon, and four-phonon infrared absorption peaks have been observed [9]. This multiphonon absorption is not strong; the absorption coefficient associated with the largest peak is about 12 cm⫺1. At higher wavelengths, the intrinsic absorption becomes negligible again.
Diamond Film Optics
613
In summary, pure diamond has an optical window spanning the ultraviolet, visible, infrared, millimeter, and microwave regions of the electromagnetic spectrum, with slight absorption in the infrared due to phonon absorption and an ultraviolet cutoff due to band-to-band optical excitation. These properties are referred to as the intrinsic optical properties of diamond, that is, those associated with carbon atoms arranged on a perfect diamond lattice. B.
Impurities and Diamond Types
Few natural single-crystal diamonds approach the intrinsic perfection. Impurities may give rise to additional electronic absorption levels and to vibration-induced absorption not found in pure diamond. Also, defects and impurities break the lattice symmetry, allowing one-phonon infrared absorption [10], which is forbidden in pure diamond. In fact, the effects of impurities on the optical properties of diamond are primarily responsible for the classification of natural diamond into four main types: Ia, Ib, IIa, and IIb [11]. Type I diamonds contain significant nitrogen impurities. Types Ia and Ib differ in that for Ib diamonds the nitrogen is present in substitutional sites and contributes to an electron paramagnetic resonance (EPR) signal. Type Ia is not EPR active, and at least in some cases the nitrogen forms aggregates or platelets within the crystal. Substitutional nitrogen gives rise to absorption in the higher energy portion of the visible spectrum, which causes the diamond to have a yellow appearance. When the nitrogen is not EPR active, the effect on color is less. For example, a type Ib diamond with 50 ppm atomic nitrogen has a pronounced yellow-gold color, whereas a type Ia diamond with 50 ppm atomic nitrogen may have only a slight tinge of color [11]. However, type Ia diamonds can have nitrogen concentrations as high as 2500 ppm atomic, and such crystals do show pronounced absorption in the visible portion of the spectrum. Type II diamond differs from type I in that type II does not contain appreciable nitrogen. Type IIa is ‘‘pure’’ diamond although purity is a matter of degree, varying from stone to stone. It has been suggested that a good specification for dividing between types I and II would be at the nitrogen level of 1 to 5 ppm atomic [11]. Type IIb is also nitrogen free but contains boron. Substitutional boron gives rise to strong infrared absorption that has a tail extending into the lower energy portion of the visible spectrum. As a result, a boron concentration of only 10 ppm produces a deep blue color in the crystal. Figure 3 illustrates comparative transmission plots for a colorless type IIa diamond, a pale yellow type Ia diamond, a gold-yellow type Ib diamond, and a blue boron-doped type IIb diamond. The type Ia in this figure is a natural stone; the others are single crystals grown synthetically
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Figure 3 The single-crystal transmission spectra of (A) a type IIa colorless diamond, (B) a blue type IIb diamond, (C) a pale yellow type Ia diamond, and (D) a golden yellow type Ib diamond [11].
by the high pressure General Electric process [11]. Of note for electronic purposes, substitutional boron acts as an acceptor, introducing holes in the valence band such that type IIb diamond is semiconducting. All four types of diamond occur naturally, with the greatest preponderance being type Ia. In one study of natural diamond, 98% were type Ia, 0.1% were type Ib, and the remainder were type II.
Diamond Film Optics
Figure 3
615
Continued
Because nitrogen is so commonly found as an impurity, most natural diamonds exhibit various shades of yellow and brown. Colorless diamonds are comparatively rare. The largest by far colorless diamond discovered to date is the Cullinan diamond found in South Africa in 1905 and weighing 3106 carats as a rough stone. It was subsequently cut into smaller stones, the largest being 530 carats. Large type IIb diamonds are quite rare, the largest being the Hope diamond, which is approximately 45 carats. Although there are several much larger diamonds, this stone, which originated in India,
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is famous for its color, a deep sapphire blue.* Also rare, but occasionally found, are diamonds of other colors including orange, pink, mauve, green, red, and black. These also result from impurities or defects. For example, manganese is associated with a broad absorption at 0.55 m that produces a pink or mauve color. Color may also be induced synthetically by creating optical centers using radiation-induced damage. Lattice damage caused by electron, neutron, and gamma ray irradiation produces an absorption band centered near 2 eV in the visible and also an ultraviolet band extending from approximately 3 eV to the absorption edge. The net effect of irradiation in a type IIa diamond is a change in appearance from colorless to green, bluegreen, or blue, depending on the irradiation [11]. Further information on color centers in diamond may be found in Chapter 3 of this book and in Refs. 4 and 5 of this chapter.
III.
OPTICAL PROPERTIES OF POLYCRYSTALLINE CVD DIAMOND
A.
Scattering Effects in CVD Diamond
The very best CVD polycrystalline diamond optical samples do in fact exhibit many optical properties quite comparable to those of type IIa singlecrystal diamond [12]. For example, in a review report by Harris [13] that summarizes the results of a multigroup effort funded by the U.S. Navy to produce optical quality CVD diamond, clear (from the band gap cutoff in the ultraviolet to the phonon absorption in the infrared) windows with thicknesses of 0.3 to 1.0 mm and diameters† up to 25 mm were described. However, much CVD diamond does not have this level of optical quality. Freestanding samples in the as-grown state are often opaque and show a color ranging from a gray to a black appearance. One factor that limits the optical performance of polycrystalline diamond films is scattering at the rough growth surface of the polycrystalline films. First, however, consider the case where scattering effects are neglected. For wavelengths below the band gap and not in the portion of the infrared where absorption occurs, a freestanding, optically smooth diamond sample should show the transmission spectra of a transparent plate in air. For such *One of several diamonds of historical interest, this gem was acquired by the royal family in France in the 17th century, disappeared in the turmoil of the French Revolution in the 18th century, appeared again in England in the 19th century, and in the 20th century became part of the Smithsonian collection in Washington D.C. † In subsequent years, the achievable diameter for high-quality optical samples of CVD diamond increased significantly.
Diamond Film Optics
617
cases, assuming that the incident light is a transverse electromagnetic (TEM) wave (which is valid except for cases of rapidly diverging or converging light), the optical power transmission for normal incidence on a diamond plate with front and back air interfaces is given by the expression [14] T=
(1 ⫺ r2ad)(1 ⫺ r2da) 1 ⫹ r2adr2da ⫹ 2radrda cos(2␦)
(8)
where the Fresnel coefficients are rad =
na ⫺ n d na ⫹ n d
(9)
rda =
nd ⫺ n a na ⫹ n d
(10)
and
at the top and bottom interface, respectively, and
␦=
2 ndt
(11)
is the change in phase of the beam on traversing the diamond film of thickness t. The constructive and destructive interference resulting from multiple reflections at the two interfaces results in a transmission oscillating between 100% and approximately 50% with peaks separated in wavenumber by integer multiples of (ndt/2)⫺1. This is illustrated in Fig. 4A, in which T is plotted as a function of wavenumber using Eq. (8). Also shown for reference is T plotted as a function of in Fig. 4B. Such oscillatory behavior can be observed experimentally throughout the spectrum if scattering is eliminated by polishing the diamond surface to optical smoothness. It is also observed on nonpolished samples if the optical wavelength is much larger than the surface roughness [15,16]. For example, Fig. 5 shows the transmission spectrum for a freestanding 76.2-m-thick film with grain size in the range 2– 8 m. The measured transmission [16] is shown as a function of wavenumber, from 20 to 350 cm⫺1, corresponding to wavelengths from 500 to approximately 29 m. In the longer wavelength portion of the transmission spectrum, the transmission behavior approximates the ideal case of Fig. 4. However, for shorter wavelengths, the transmission drops off because of scattering at the rough surface. Scattering at a rough surface affects both transmitted and reflected light. For light entering a rough diamond surface from air, the roughness of the surface causes phase differences in transmitted light. This introduces a correction factor for power transmission which, following Filinski [17], may be written as
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Figure 4 Ideal transmission for a 75-m-thick freestanding, optically smooth diamond plate. (A) The regularly spaced peaks when plotted as a function of wavenumber; (B) the same information plotted as a function of wavelength.
Diamond Film Optics
619
Figure 5 Measured transmission for a 76.2-m-thick film with grain size in the range of 2 to 8 m. (From Ref. 16.)
ST = exp
冋冉
冊册
2 (nd ⫺ na ⫺
2
(12)
Here is the free-space light wavelength and is the root mean square (rms) of the surface roughness, assuming a Gaussian distribution of surface height about the mean and small compared with the wavelength [18]. Additional models have been developed for scattering of shorter wavelengths from rough surfaces [19]. The rough surface also affects the interference term due to multiple reflections within the diamond film. The internal power reflection at a rough surface is multiplied by a correction term [17] SR = exp
冋 冉 冊册 4 nd ⫺
2
(13)
Again, the reduction of intensity is caused by the phase fluctuations appearing in the reflected beam. Generally on CVD diamond films, the top surface, or growth surface, is much rougher than the bottom surface. Considering only the top surface to be rough and assuming that the electric field reflection coefficient is corrected by a factor (SR)1/2, then the expression for T is modified to account for the rough surface to become TR, where [20] TR =
(1 ⫺ r2ad)(1 ⫺ r2da) ST 1 ⫹ SRr2adr2da ⫹ 2radrda兹SR cos(2␦)
(14)
Figure 6 shows the theoretical transmission for a 75-m-thick freestanding plate with an rms surface roughness of 2 m on the top surface. Comparing Fig. 6 with Fig. 4, significant loss in transmission is observed in both the
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Figure 6 Theoretical transmission for a 75-m-thick freestanding diamond film that has an rms surface roughness of 2 m on one side and is optically smooth on the other side. Results may be compared with the theoretical transmission for a smooth diamond film in Fig. 4 and the experimental results of Fig. 5.
experimental and theoretical cases as the wavelength decreases. For in the visible portion of the spectrum, the transmission would be quite small in both cases. Partly as a result of scattering, as-grown polycrystalline CVD diamond samples often have both reflectance and transmittance values less than 1%, far different from the high values associated with gem quality single-crystal diamonds. The resulting optical appearance has caused more than one diamond researcher to be faced with incredulity when presenting CVD samples to the public for inspection. The general conclusion to be drawn from the literature regarding optical transmission through, and reflection from, polycrystalline diamond is that both are often limited by the surface roughness of the growth surface [10,21–24]. In the far infrared, where the wavelength is large compared with surface roughness, the transmittance approximates that of a transparent plate in air. At smaller wavelengths the transmittance drops off due to scattering at the rough surface. The smoother the surface, the shorter the wavelength at which appreciable transmittance can occur. Consequently, surface roughness can be a major impediment for applications. Methods for reducing
Diamond Film Optics
621
this problem by growing smooth surfaces and by postprocessing are addressed in Secs. IV and V of this chapter. B.
Impurity-Induced Absorption in CVD Diamond
CVD diamond films often contain chemical impurities that decrease optical transparency in portions of the optical spectrum. Some of these impurities are also found in natural diamond and play optical roles similar to those in single-crystal diamond. Examples of these are nitrogen (due, for example, to air contamination in the CVD vacuum system or in the feed gases) and boron (intentionally introduced as a p-type dopant). Nitrogen introduces states in the energy gap reported variously to be between approximately 1.6 and 1.9 eV below the conduction band. These states cause absorption at photon energies below the intrinsic energy gap of diamond, in the ultraviolet and blue portion of the spectrum, such that a yellow coloration results. Boron introduces states in the gap at approximately 0.4 eV above the valance band. These states cause absorption that peaks at approximately 2.5 m wavelength but tails to approximately 0.5 m in the visible, causing a blue coloration [10]. However, CVD diamond often has additional impurities or defects not commonly associated with natural single-crystal diamond. If the CVD system contains metallic electrodes or filaments, metallic contaminants in the film can result. Likewise, if the CVD system involves a discharge contained within silica walls, both silicon and oxygen may be incorporated in the film. Feed gases are also sources of impurities. For example, hydrogen is commonly incorporated in feed gases, both directly and as a component of the carbon-containing molecule. As a result, hydrogen is incorporated in the growing film. When the CVD chemistry involves oxygen-containing species, the CVD diamond film also contains oxygen. In addition to chemical impurities, diamond films contain structural defects. In particular, depending on growth conditions, non-sp3 carbon bonds may be present, corresponding to nondiamond carbon, giving rise to additional allowed energy levels and correspondingly additional optical transitions. Other structural defects include crystallographic defects (twins, stacking faults, and dislocations) as well as grain boundary regions between crystallites. A variety of analytical methods have been used to assess defect and impurity concentrations in diamond films including nuclear magnetic resonance (NMR) [25], x-ray photo-electron spectroscopy (XPS) and x-ray diffraction [26], electron paramagnetic resonance (EPR) [27], photothermal deflection spectroscopy (PDS) [28], transmission electron spectroscopy (TEM) [29,30], cathodoluminescence [31,32], Raman spectroscopy [33,34], Auger electron spectroscopy, secondary ion mass spectrometry (SIMS), and electron energy loss spectroscopy (EELS) [35]. An example of an impurity anal-
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ysis of CVD diamond by McNamara et al. [26] with regard to hydrogen, oxygen and nitrogen is shown in Table 1. Considering the nitrogen content, it is noted that amounts in this table range from 20 to 100 ppm, which, following the discussion in Sec. II, corresponds to type I diamond. Additional information concerning the characterization of impurities and defects in diamond films may be found in Chap. 3 of this book. Various impurities cause absorption beyond that of the intrinsic optical absorption of diamond by allowing either additional photon-induced electronic transitions or additional photon-induced lattice vibrations. In the case of hydrogen it is principally the latter, such that the effect of carbon-hydrogen bonds in diamond is to cause additional absorption in the infra-red. This added absorption is particularly pronounced in the CH stretch region (2750– 3050 cm⫺1 wavenumbers or approximately 3.5 m wavelength); however, additional absorption due to hydrogen incorporation is also observed in other parts of the infrared spectrum. In the study reported by Harris [13], which included diamond samples grown by various methods including hot-filament, microwave plasma, and dc arcjet CVD processes, hydrogen contents ranged from 16 ppm to fractions of an atomic percent depending on the growth method and gas feed composition. At the higher hydrogen concentrations, C — H stretching absorption bands dominated the infrared absorption spectra with a strong absorption peak near 3.5 m. At the lower range of hydrogen content in the film, the effect of hydrogen on absorption is negligible. Figure 7 illustrates infrared absorption spectra in hydrogen-contaminated polycrystalline diamond [11]. Not only is additional optical absorption
Table 1 Example of Impurity Content Analysis of Diamond Films Deposited by Microwave Plasma-Assisted CVDa Sample A B C D E F G H J a
Hydrogen (%)
Nitrogen (%)
Oxygen (%)
0.085 0.180 0.027 0.045 0.021 0.043 0.020 0.015 0.092
0.0048 0.0047 0.0020 0.0035 0.0018 0.0088 0.0024 0.0039 0.0101
0.0091 0.0075 0.0053 0.0157 0.0034 0.0832 0.0132 0.0032 0.1970
Results are expressed in atomic percent. Source: Ref. 26.
Diamond Film Optics
623
Figure 7 Absorption spectra showing the effect of hydrogen impurities in polycrystalline diamond. The impurities activate one-phonon absorption and also cause absorption due to C — H bonds [11].
present due to C — Hx stretching modes, but also the presence of hydrogen impurities activates one-phonon absorption, which is forbidden in pure diamond. Consequently, samples with strong C — Hx absorption also exhibit strong one-phonon absorption. This single-phonon absorption is also illustrated in Fig. 7 in the form of absorption in the range 8–10 m. Most materials transparent in the region 8–10 m are soft, making diamond an attractive alternative for infrared applications. To the extent, therefore, that hydrogen is incorporated into CVD diamond, this issue can be a performance-limiting concern for CVD diamond. Based on the results shown in Fig. 7, if the hydrogen content is reduced to less than 0.02% atomic, such absorption is negligible even for thick films. For thinner films the effect on percent transmission will, of course, be correspondingly less. However, experimental observation of infrared absorption due to hydrogen impurities
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has also been reported on thin freestanding membranes [15,36]. For example, Johnson and Weimer [15] removed the silicon substrate from diamond films by etching to study the transmission of a 3-m-thick film and observed small but detectable impurity absorption peaks at 2925–2830 cm⫺1 and 3325 cm⫺1, which were attributed to C — Hx groups. Significantly larger peaks were observed on a 13-m-thick film. Fortunately, if appropriate means are taken in choosing deposition parameters, the hydrogen content of CVD diamond can be reduced to a point at which the effect on absorption is negligible. Unfortunately, this generally involves the use of relatively low methane concentrations, which results in a trade-off between optical quality and growth rate. The other elements in Table 1, oxygen and nitrogen, also play major roles in the optical properties of CVD diamond. The effect of nitrogen with regard to causing increased absorption in the near UV and visible portion of the spectrum has been previously noted. EPR data suggest that most nitrogen in CVD diamond is present as point-defect substitutional impurities rather than aggregates [27], which in terms of traditional diamond classification corresponds to type Ib rather than type Ia. In terms of oxygen, this element is sometimes added to the gas flow directly as O2, or in the form of CO or CO2, in order to reduce the graphitic content of the diamond film and thereby improve optical transmission in the visible [37]. However, both oxygen and nitrogen also cause increased absorption in the infrared. IR spectroscopy has identified absorption due to a variety of oxygen- and nitrogen-containing groups, including O — CH3 and N — CH3 [26] which introduce absorption in the 8–12 m wavelength region. Therefore, adding oxygen in order to improve visible transmission may be accompanied by a trade-off in terms of decreased transmission in the far infrared. Also noted is that even small leaks in CVD diamond systems can lead to observable oxygen and nitrogen incorporation. In addition to hydrogen, oxygen, and nitrogen, optical effects due to other chemical impurities have been reported. Lithium, of interest as an ntype dopant, has been reported to cause increased absorption in the vicinity of 1.5 eV [28]. The incorporation of silicon has been reported to be associated with absorption at 1.68 eV [31]. C.
Defect-Induced Absorption in CVD Diamond
Structural defects other than chemical impurities are also strongly correlated with optical properties of CVD diamond. Principal among these is the effect of graphitic carbon or sp2 bonds. An early and instructive study in this regard was reported by Collins et al. [31] in which absorption and cathodoluminescence spectra of polycrystalline diamond were investigated for different
Diamond Film Optics
625
growth conditions. The diamond films were deposited on silicon using microwave plasma-assisted CVD systems and methane-hydrogen feed gas mixtures with methane concentrations varying from 0.3 to 3.0%. In order to avoid scattering effects on the optical absorption measurements, the rough growth surface of the diamond was mechanically polished. The central region of the silicon wafer was then etched away, leaving a supported diamond film that was smooth on both sides. In addition to the CVD films, isolated CVD crystals were investigated by growing individual crystals about 100 m across on sparsely seeded substrates. These individual crystals were detached from the substrate and mounted on optical sample holders. As shown in Fig. 8, the results of Collins et al. [31] showed that absorption in the films was a strong function of methane concentration and of a type not attributable simply to C — H bonds. All samples exhibited an almost monotonically increasing absorption with increasing photon energy. The band gap absorption edge is clearly visible in Fig. 8 for the 0.3% methane film; however, sub–band gap absorption is appreciable with ␣
Figure 8 Absorption spectra of CVD diamond films grown with varying methane concentrations as noted, as well as the absorption spectrum of an isolated CVD crystal. (From Ref. 31.)
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reaching approximately 500 cm⫺1 prior to the onset of band-to-band generation. The color of this film was described as pale yellow, consistent with appreciable absorption in the higher energy portion of the visible. For the higher methane concentrations, the film samples were too strongly absorbing to extend the measurements into the ultraviolet region. The 1.5% methane film shows an absorption coefficient of over 2000 cm⫺1 in the blue portion of visible spectrum and exhibited a dark brown color to the eye. Based on Raman analysis, evidence suggested that the absorption evident in Fig. 8 was due to nondiamond carbon in the form of sp2 bonds. The importance of nondiamond carbon bonds in causing appreciable sub–band gap absorption was further confirmed and quantified in later studies. In an investigation by Rosa et al. [38] of microwave plasma-assisted CVD diamond, a set of diamond films were deposited on silicon under variable growth conditions such that films ranged from white polycrystalline films to brown nanocrystalline films. The value of ␣ at 2.41 eV ranged from a low of 3 cm⫺1 for the least absorbing films to as high as 500 cm⫺1. Films were also characterized by a Raman quality factor defined as
=
ID ID ⫹ IN
(15)
where ID and IN are the integrated intensities of the diamond and nondiamond Raman signals, respectively. From , a volume fraction of the nondiamond component was calculated and correlated with the absorption coefficient at 2.41 eV with an approximately one-to-one correlation. When the absorption coefficient is nonnegligible, the expression for transmission is modified as shown in Eq. (16), which accounts for both scattering due to a rough surface and absorption within the film to arrive at a transmission, TRA, given by [22,39] TRA =
(1 ⫺ r2ad)(1 ⫺ r2da)exp(⫺␣t) 1 ⫹ S r r exp(⫺2␣t) ⫹ 2radrda兹SR exp(⫺␣t)cos(2␦) 2 2 R ad da
(16)
The importance of absorption clearly depends on the diamond thickness. For an absorption coefficient of 500 cm⫺1 and a 1-m-thick film, the transmission loss resulting from absorption is approximately 5%, which may be negligibly small for many applications. However, for a 500-m-thick freestanding diamond sample, the transmission loss would be 99.3%. Because the optical properties of CVD diamond samples may depend on both wavelength-dependent scattering and various wavelength-dependent absorption mechanisms, separation of optical constants is nontrivial. The procedure involves best-fit solutions to the optical equations using floating values of the free parameters [22,39–41].
Diamond Film Optics
IV.
FREESTANDING DIAMOND AND DIAMOND DIAPHRAGMS
A.
Application and Growth Considerations
627
One class of optical applications of diamond involves the use of either freestanding diamond or diamond diaphragms as optical elements. This section discusses a variety of application-related issues associated with such structures. Various applications for diamond windows have been reviewed by Seal and van Enckevort [1], most of which involve operation in difficult environments. Diamond windows have been used in spectroscopic instrumentation for various spacecraft including the Nimbus and aforementioned Pioneer series as well as the Galileo Jupiter probe. On earth, diamond windows are used in the chemical industry for spectroscopic quality control of various caustic, high-temperature processes and also in laboratories. Because of diamond’s resistance to scratching and ease of cleaning, diamond components are used in instrumentation by the food industry for contact analysis by optical means. In the field of medicine, diamond optical elements have found use as miniature endoscope windows. The transmissive properties of diamond have also been employed for medical use in diamond blades (which have superb cutting edges) in that laser energy can be piped to the cutting edge, where it is used to cauterize blood vessels. In some cases, laser energy at the blade edge also seems to facilitate the cutting of certain tissues. Freestanding diamond has also been used for high power light transmission windows [42,43]. The most common optical material, silica-based glass, is not well suited for high-power windows because its thermal conductivity is so low, typically on the order of 0.01 W/cm deg. The thermal conductivity of diamond is over a thousand times higher and the optical absorption coefficient of diamond can be as low as or lower than that of glass. Consequently, the optical power transmission of diamond is considered to be at least 1000 times higher than that of glass. This is relevant to the designers of high-power lasers, where limits on power are set by the damage limits to laser optics rather than limitations of the laser medium or pumping mechanisms. It is also relevant to high-power gyrotrons in the mm wave portion of the spectrum. A major limitation to optical diamond applications has historically been size limitation imposed by cost. Whereas the cost of a 0.5-mm-diameter window of single-crystal diamond is measured in tens of dollars, increasing the diameter by an order of magnitude to 5 mm causes the price to be measured in thousands of dollars. To increase the diameter another order of magnitude to 50 mm is essentially to leave the realm of availability of singlecrystal diamond. It is for this reason that CVD capabilities offer the potential to greatly increase the applicability of diamond to optics. In fact, by the
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Figure 9 Measured transmission for a polished, 350-m-thick film (solid line) and a 275-m-thick type IIa crystal (dashed line). The data are smoothed so as not to show the interference oscillations. Evident in both cases are the intrinsic multiphonon absorption in the infrared between 4 and 7 m and the onset of band gap absorption in the UV [13].
early 1990s, the ability to grow freestanding diamond of significant size with transmission characteristics throughout the entire optical range comparable to intrinsic diamond had been demonstrated. Figure 9 shows transmission results for a 350-m-thick CVD diamond sample, fabricated at Raytheon in 1992 [13]. Data are plotted in direct comparison with transmission results for a 275-m-thick type IIa single crystal, and essentially identical properties are observed.* Evident in both samples is the onset of band gap absorption at 0.23 mm and the multiphonon absorption inherent to the diamond lattice in the infrared. Absent are any significant effects due to defects and impurities as discussed in Sec. III of this chapter. Within 2 years, such *The data in Fig. 9 are smoothed so that interference effects are not seen. Also, the CVD sample was polished to optical smoothness after growth.
Diamond Film Optics
629
visually transparent samples were available from the same company at 60 mm diameter [44] using microwave plasma-assisted CVD. Gray et al. [45,46] at Norton Co. reported 100-mm diameter CVD ‘‘white diamond’’ wafers, 1 mm thick, with infrared transmission and Raman characteristics comparable to those of type IIa diamond in approximately the same time frame, 1992. In fact a variety of CVD methods have been used to deposit freestanding or supported diaphragm diamond for optical applications or investigations, including hot filament [47], arc plasma jet [45,46,49], and microwave plasmas [16,50–52]. Regardless of deposition method, an important practical consideration for freestanding diamond elements is that of the available deposition rate. Unsupported diamond layers intended for optical applications must be of sufficient thickness to provide necessary robustness. Generally such films are hundreds to several hundreds of micrometers thick. High deposition rates are therefore desired for reasonable production times. However, high deposition rates do not necessarily yield the best quality diamond. In the previously mentioned U.S. Navy–funded study [13], it was found for all production methods used in the investigation (hot filament, arcjet, and microwave) that maximum growth rates were limited to 1 to 5 m/hr for optical quality diamond. Higher growth rates were obtainable by increasing the methane concentration, but this approach to high deposition rates yielded lower optical quality. Therefore, one challenge for freestanding CVD optical diamond is to achieve high deposition rates while maintaining little or no sp2-bonded carbon and a low concentration of C — H bonds, thereby maintaining the low absorption coefficient necessary for high transmissivity in thick films. Issues specifically related to increasing growth rates are addressed in other chapters in this book that are related to deposition methodologies. B.
Diamond Surface Considerations
Another important consideration for applications is the diamond surface condition. As previously noted, the rough growth surface of a polycrystalline diamond film presents an important performance limitation due to optical scattering. Generally, the surface roughness of diamond films increases as the thickness increases. In the course of film growth, growth competition between the individual crystallites results in more and more crystallites being buried and a decreasing number of surviving crystallites at the surface with the size of the crystallites increasing as growth continues. Consequently the growth surface of an as-grown CVD diamond sample with a thickness of 500 m may show well-faceted crystallites with average sizes of 50 m and larger [45]. Such a surface can easily have an average surface roughness
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Ra value of several micrometers. For such samples, except for the far infrared, optical transmission losses due to scattering are appreciable and it is necessary to polish the diamond surface in order to achieve good transmission in the visible and ultraviolet. Because diamond is the hardest known material, polishing a rough diamond surface is a nontrivial task. However, a variety of methods have been developed. Drawing from experience in the gem industry, a traditional method is to rub the rough diamond surface against a cast iron surface. One version of this approach is to rotate an iron wheel (called a scaife) at high speed (2000 rpm) and press the diamond against the rotating surface.* Significant heat is developed by the friction between it and the diamond. Another approach is to heat the iron to approximately 350⬚C and use slow rotation [50]. The polishing in either case is believed to be due to a chemical reaction between the diamond carbon and the iron to form iron carbide, which is subsequently oxidized in air, due to the high temperature, and is easily removed by mechanical polishing. This method can take several weeks to achieve an optically smooth surface. If polishing is too aggressive, the diamond piece can be shattered. Alternatively, the surface can be polished with high-speed lapping using diamond grit as the abrasive [53]. Surface smoothing and diamond removal have also been achieved by using diffusional reactions with hot metal plates. In this approach, the diamond workpiece is placed in static contact with iron, manganese, cerium, lanthanum, or La-Ni eutectics at 600–900⬚C in an argon atmosphere, under which conditions diamond dissolves in the metal [54–58]. For example, as reported by Jin et al. [55], a freestanding CVD diamond piece (1 cm ⫻ 0.5 cm ⫻ 250 m thick) was sandwiched between metal sheets with compressive stress of 3000 psi for iron and 200 psi for manganese for 48 hr at 920⬚C. The rough growth facets were reported to be mostly removed and the diamond thickness was reduced by approximately 50 m. Using rare earth metals, removal rates were more than five times higher because of the large liquid solubility of carbon at high temperatures. Noncontact methods including ion beams, plasmas and laser treatment have also been investigated for improving the diamond surface [59–64]. This approach can be facilitated by first applying a sacrificial planarizing layer to the diamond surface and then using a material removal process that removes the diamond and planarizing layer at comparable rates. For example, Bovard et al. [62] have described the use of photoresist/titanium-silica as a planarizing layer followed by reactive ion beam etching with 500 eV
*Iron is the most common scaife material but not the only possibility. Nickel, for example, can also be used.
Diamond Film Optics
631
oxygen ions to achieve optical quality smoothness. Starting with a 5-cmdiameter CVD sample with a 1-m peak-to-valley surface roughness, the reactive ion beam polishing procedure reduced the peak-to-valley distance to 55 nm and resulted in an rms surface roughness of 5.5 nm [62]. In some cases freestanding diamond is grown to a thickness greater than that needed for the final product so that postgrowth diamond removal can be used to correct for nonuniform thickness or sample bowing. Noncontact methods, such as electron-cyclotron-resonance (ECR) oxygen plasmas, have been shown to be much faster than conventional scaife or lapping methods for bulk removal of diamond from workpieces [65]. For example, Chakraborty et al. [66] have reported the use of ECR plasmas to etch 10cm-diameter CVD freestanding diamond, 1 mm thick, with O2/SF6/Ar plasmas. Diamond removal rates of 7 m/hr with ⫾10% uniformity across the substrates were reported. The etch mechanism in such cases is primarily highly anisotropic ion-assisted reactive etching by diamond. To first order, the etch rate increases linearly with increasing ion energy. Therefore, high etch rates are facilitated by substrate bias. The role of SF6 is to avoid the formation of a residual black film on the diamond surface [66,67]. Such etching erodes the rough crystal facets, but it does not produce an optically smooth finish, which requires further processing. Although it is possible to postprocess CVD diamond surfaces to optically smooth surfaces by various polishing means, it is clearly of advantage to grow the films with smooth surfaces such that the need for polishing is greatly reduced or eliminated. One approach, described by Wild et al. [51,68,69], is to grow textured diamond films on silicon substrates with a preferred orientation of directional growth such that a relatively flat growth surface is covered with {100} faces results. After growing a thick diamond film, the silicon substrate is removed, leaving a freestanding diamond sample. By this growth method, 150-m-thick films were deposited with surface roughness values of 150 nm. This Ra value is still an order of magnitude higher than that needed for lossless transmission throughout the entire visible spectrum, but it is substantially less than for polycrystalline films grown by conventional means. Such orientation is achieved by adjusting growth chemistry such that the ratio of growth rates on the {100} and {111} faces is slightly below 兹3. In this case, cubo-octahedral nuclei with small {100} faces develop. The direction of fastest growth is perpendicular to these faces, and as the film grows thicker the sizes of these {100} faces increases. Consequently, growth techniques that result in preferred growth orientations are useful for optical applications. Table 2 summarizes a specific method as described by Maeda et al. [70] for growing diamond that is textured in the 具100典 direction and results in highly oriented {100} faces on silicon substrates in a microwave plasma. This approach uses a three-step
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Table 2 Three-Step Microwave Plasma CVD Conditions for Growing Highly Oriented {100} Diamond Films on Silicon Parameters CH4 (sccm) H2 (sccm) Bias voltage (V) Treatment time (min) Microwave power (W) Substrate temperature (⬚C) Pressure (kPa)
Carburization 2 98 0 15–120 570 830 2.65
Bias treatment 2 98 ⫺150 1–15 830 830 2.65
Growth 0.5–2 98–100 0 >120 460–670 800–900 4
Source: Ref. 70.
growth sequence in which the first step carburizes the surface, the second step adds an electrical bias to the substrate which creates a textured growth surface, and the third step involves the growth of the bulk of the film. Film texture has also been correlated with hydrogen incorporation and therefore with infrared absorption. Haq et al. [71] investigated infrared absorption for various film textures and observed that the lowest C — H absorption was associated with 具110典-oriented films, followed by 具100典 and then 具111典. Single-phonon absorption was also found to be highest in the case of 具111典 textured films and lowest for 具110典 and 具100典 with no significant difference between the latter two. An alternative approach used to achieve smooth surfaces that works for relative thin films is to deposit very fine grain films [69,72,73]. If the nucleation density is very high, on the order of 1010 cm⫺2 or higher, and the film thickness is kept sufficiently small that large grains do not have a chance to form, then quite small surface roughness values can be achieved. Yang et al. [74] reported nucleation densities as high as 1011 cm⫺2 and a mean surface roughness of 30 nm for 1-m-thick films. The role of fine grain films is discussed further in this chapter in the context of substrate coatings in Sec. V.B. C.
Mechanical Considerations for Plates and Diaphragms
For several optical applications of freestanding diamond, mechanical strength is an important figure of merit as well as optical transparency. Whereas it is possible to make freestanding diamond pieces of significant size that essentially have the optical transmission of intrinsic diamond, the
Diamond Film Optics
633
same cannot be said of mechanical strength. The mechanical strength of optical quality 0.5- to 1-mm-thick CVD diamond has been reported to be significantly less than that of natural diamond [53]. Optical quality diamond has been reported to grow in a highly stressed state that can lead to macroscopic and microscopic cracks [13].* It has been suggested that the incorporation of voids and defects arises as a result of random in-plane grain orientations coupled with the overgrowth of twin boundaries and that this is relatively independent of reactor design [75]. Such cracks and defects play a critical role in mechanical strength because the strength of polycrystalline CVD diamond is governed by flaws. One way to measure mechanical strength is by means of the ring-onring test, in which one side of a diamond plate is supported on a metal ring and the other side is contacted by a metallic load ring of smaller radius. The load magnitude is gradually increased until the point of fracture. In one study of ring-on-ring load testing of CVD diamond plates produced by a variety of deposition methods, measured strengths were reported to be an order of magnitude less than would have been expected for single-crystal diamond [13]. The fracture origins in the CVD diamond plates corresponded to relatively large flaws, on the order of 100–300 m in size for 800-m-thick disks. In a study by Valentine et al. [76], the bursting strength of diamond plates with thicknesses in the range of 179–319 m and diameters in the range of 14–25 mm were measured by applying gas pressure to diamond disks supported at the disk periphery. In this case, the CVD diamond strength was reported to be about half that of natural diamond. Fine polishing of the surface seems to have no effect on strength [13]. For some optically related applications, it is possible to consider the use of diamond membranes, or diaphragms. For these structures the mechanical strength of CVD diamond does not necessarily differ significantly from that of single-crystal diamond, perhaps because the thinner films cannot have flaws as large as those found in the thicker structures [53]. Diamond diaphragms subject to loads can, however, show considerable deflections, resulting in distorted optical surfaces. The load deflection behavior of pointloaded CVD diamond thin-film diaphragms was investigated by Glime et al. [77]; circular diaphragms consisting of diamond film 1.5 m thick and 2– 4 mm in diameter adhering to a silicon rim were prepared by mechanically grinding and chemically etching the single-crystal substrate. A nanoidentor was used to provide load versus deflection data. The maximum deflections
*It should be noted that natural diamonds are also occasionally found that are in highly stressed states. There are reports of large stones of considerable potential value spontaneously shattering when brought to the surface of the mine because of temperature changes.
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were observed to be up to five times the film thickness before failure. When the central deflection of a plate exceeds about one half the plate thickness, classical plate theory relating load and deflection to the mechanical material properties is no longer valid and nonlinear models must be used [78]. The relationship between applied central point load, P, deflection at the center of the diaphragm, w, and diamond thickness, t, is reported to be of the form 4Et4 P= 3(1 ⫺ 2)r2
冋
w ⫹A t
冉 冊册 w t
3
(17)
where r is the radius, E is Young’s modulus, is Poisson’s ratio, and A is a constant equal to about 0.43 for a point-loaded circular plate loaded in the nonlinear region and diamond diaphragm deflection data agrees well with this equation [77,79]. Diamond diaphragm structures often contain internal stresses. Berry et al. [80] have used vibrating membrane methods to measure internal stresses in CVD diamond deposited on silicon using microwave plasma deposition. Central regions of the substrates were back etched to leave the diamond film as a taut membrane, with net tensile stress, supported circularly at the edges. Film stresses ranging from 9.4 to 139 MPa were measured, depending on growth temperature. Flowers et al. [81] found the burst pressure for diamond diaphragms to increase linearly as a function of a parameter R, where R is the square root of the film thickness divided by the diaphragm diameter. In some cases, compressive internal stresses in diamond diaphragms have resulted in film buckling, causing wrinkling and nonplanar surfaces [77]. The issue of internal stresses in diamond films is discussed farther in the context of film coatings in Sec. V.C. Mechanical strength is also of interest for shapes more complex than simple plates. Of particular interest for certain airborne applications has been the fabrication of freestanding CVD diamond domes such as could be used to protect nose cone infrared sensing instrumentation where aerodynamic considerations are crucial [53,82]. Near net shape domes have been reported by depositing on nonplanar growth surfaces. Potential applications of such window structures involve speeds as high as mach 4.* Consequently, significant differential pressures and complex strain patterns must be considered [83]. Another mechanical strength–related figure of merit particularly relevant for airborne applications is the erosion resistance to liquid impact. In
*The limit on speed is imposed by temperature considerations. At mach 6, stagnation temperatures are such that diamond may convert to graphite.
Diamond Film Optics
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a study by Jilbert et al. [84], the rain erosion resistance of diamond was investigated by using high-velocity jets, 0.8 mm in diameter, designed to simulate the effects of spherical raindrop impact. Using 1-mm-thick samples, the threshold velocity for catastrophic failure was evaluated for both type IIa natural single crystals and CVD diamond. For the type IIa samples, the threshold velocity for catastrophic failure was 515 m/sec and for CVD diamond the threshold velocity for one sample was 325 m/sec and for another was 375 m/sec. For sand erosion, the CVD diamond was again found to be considerably weaker than its natural counterpart, as illustrated in Fig. 10. Although CVD diamond is inferior to single-crystal diamond in regard to mechanical strength, it should be noted that competing infrared-transmitting windows all suffer to some degree from resistance to erosion and other issues related to mechanical strength. A bursting strength that is half that of single-crystal diamond is still approximately an order of magnitude higher than those of other materials transparent at an IR wavelength of 10.6 m, such as GaAs, Ge, and ZnS [76]. Also, figures of merit for a given application generally involve multiple material parameters that can further favor diamond. For example, CVD diamonds offer a useful combination of optical transmission and resistance to erosion and thermal shock. Harris [53] noted that even with a relatively low mechanical strength, the extremely high ther-
Figure 10 The exposure time required to form cracks in natural and CVD diamond upon exposure to sand erosion with a flux of 10.5 kg m⫺2 sec⫺1. (From Ref. 84.)
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mal conductivity of CVD diamond coupled with its low thermal expansion provides a thermal shock resistance that is two orders of magnitude greater than that of other optical window materials. Many other aspects of CVD diamond windows compare well with single-crystal diamond. It is noted that the hardness and modulus of CVD diamond are found equivalent to those of type IIa natural diamond [53]. Also, high optical quality and high thermal conductivity go hand in hand for CVD diamond. In fact, optical absorption measurements have been proposed as an inexpensive technique for estimating the thermal conductivity of CVD diamond [85]. Optical quality CVD diamond thermal conductivity is reported to have room temperature conductivities of approximately 2000 W/m deg. Combined with low absorption over most of the optical spectrum and a low thermal expansion coefficient, 1 ppm/K at room temperature, the same as for type IIa diamond, CVD diamond is capable of providing the attributes required for high-power-transmitting windows.
D.
Effects of Oxidizing Environments
Finally, whether diamond is in the form of single-crystal or CVD polycrystalline, unprotected diamond is not suitable for use at high temperatures in oxidizing environments. Diamond maintains its unique transmissivity at high temperatures in nonoxidizing atmospheres and the refractive index has been measured up to at least 1200⬚C [86]. However, it is oxidized in air at temperatures above approximately 700⬚C. The resilience of polycrystalline CVD diamond varies considerably with film quality. Bachmann et al. [87] noted that some samples survive 20 min in oxygen at 700⬚C while others disappear completely within only 5 min [87]. Raman and SEM analysis indicated that defective diamond and nondiamond carbon composites are preferentially etched. Nitrogen-doped films and films grown in deposition systems with air leaks showed poor resistance to oxygen etching because nitrogen induces the decoration of diamond grains with highly defective or nondiamond carbon. The operating temperature range of diamond in oxidizing environments can be increased by applying protective coatings. An added advantage of such coatings would be to act as an antireflection layer. Moran et al. [88] have considered a variety of materials for high-temperature coatings including selected oxides, hydrides, carbides, fluorides, and nitrides of yttrium, hafnium, cerium, aluminum, silicon, and boron. In nonoxidizing environments, an upper limit for diamond applications would be 1600⬚C, at which diamond transforms to graphite, which is the stable phase of carbon at atmospheric pressure [13].
Diamond Film Optics
V.
DIAMOND-COATED OPTICAL SUBSTRATES
A.
Optical Considerations
637
The attractive surface properties of diamond—extreme hardness and chemical resistance—can be transferred to optical substrates that lack these features by means of thin-film diamond coatings. This section considers the optical properties of a system consisting of a diamond film on a transparent substrate. Consider, for example, a diamond film of thickness t on a glass substrate as shown in Fig. 11. For the ideal case of zero absorption in the diamond and of smooth surfaces such that there is no scattering, the transmission is limited only by specular transmission at the interfaces. Considering light again in the form of a TEM electromagnetic wave incident on the diamond, the light in the glass substrate is determined by a modification of Eq. (8) as TG =
(1 ⫺ r2ad)(1 ⫺ r2dg) 2 2 1 ⫹ rad rdg ⫹ 2radrdg cos(2␦)
(18)
where the coefficients rad and ␦ are given in Eqs. (9) and (11) and rdg =
nd ⫺ n g nd ⫹ n g
(19)
Equation (18) does not account for reflection at the glass-air interface or for multiple reflections in the glass. The observed transmission for the total structure, defined as TT, is related to TG as
Figure 11
Considerations for the ideal transmission of a diamond film on glass.
638
Reinhard 2 TT = TG(1 ⫺ rga )[1 ⫹ r2ga(1 ⫺ TG)] ⫹ higher orders
(20)
where rga =
ng ⫺ n a ng ⫹ n a
(21)
The higher order terms typically contribute on the order of 0.1% and may be neglected for most purposes. Equation (20) is based on the assumption that the light reflected within the glass substrate is not sufficiently monochromatic to produce interference effects over the width of the glass substrate. Consider, for example, the case where the glass is 1 mm thick and the incident light is from a monochromator with a bandwidth of 10 nm. For these conditions the wavelength separation of maximum in transmission due to multiple reflections within the glass is much smaller than the bandwidth of the monochromator—the light behaves incoherently as far as multiple reflections in the glass are concerned. Figure 12 shows the ideal transmission resulting from a 1-m-thick diamond film on glass using Eq. (20) and refractive index values of 1.54 and unity for glass and air, respectively, where Eq. (3) is used to generate refractive index values for diamond. Note that the glass, whose refractive
Figure 12 Calculated transmission for a diamond film on glass with zero scattering and zero absorption.
Diamond Film Optics
639
index is between those of diamond and air, acts as an antireflection layer, increasing the average power transmission. For a diamond plate with air interfaces on both sides, the average power transmission is given by the average of Eq. (8), which is approximately 67% for a diamond refractive index of 2.4. For a diamond-glass combination with air interfaces on top and bottom, the average power transmission is given by the average of Eq. (20), which is approximately 75%. For semiconductor infrared detector substrates, a diamond overlayer acts as an antireflection coating. For example, for a semiconductor with a refractive index of 3.6, a diamond coating would increase the fraction of light entering the semiconductor from 68 to 79%. As has been previously noted, a rough top surface will cause scattering losses due to phase differences in both transmitted light and internally reflected light. Following the treatment of Sec. III.A, the light entering the glass substrate under conditions of a rough diamond surface, defined here as TGR, is given by a modification of Eq. (14) such that TGR =
(1 ⫺ r2ad)(1 ⫺ r2dg)ST 1 ⫹ SRr2adr2dg ⫹ 2radrdg兹SR cos(2␦)
(22)
and the expression for transmission through the entire structure, accounting for reflection loss at the glass-air interface, defined here as TTR is given by TTR = TGR(1 ⫺ r2ga)(1 ⫹ r2ga(1 ⫺ TGR)) ⫹ higher orders
(23)
Figure 13 illustrates the calculated transmission using Eq. (23) for rms surface roughness of 40, 25, and 10 nm throughout the same visible/nearinfrared spectral region used to illustrate the ideal (no scatter) case of Fig. 12. For the case of 40-nm surface roughness, the effect on scattering in the visible range is nonnegligible, with the transmission value dropping to less than 40% in the blue portion of the spectrum. However, for the case of 10nm surface roughness the transmission compares very closely to the results of Fig. 12 for the ideal film without scatter. B.
Diamond-Coated Substrate Examples
One approach to applying diamond to an optical substrate is to form the two materials separately and then bond them together. For example, Partlow et al. [89] have described an ‘‘optical brazing’’ process that results in the thermal bonding of diamond to infrared-transmitting windows of ZnSe and ZnS. In this method, diamond is deposited on silicon by microwave plasmaassisted CVD. A chalcogenide bonding glass, whose refractive index is closely matched to diamond, is used as an interlayer between the diamond and, for example, ZnSe. A sandwich is made of the II-VI compound, the chalcogenide, and the diamond-coated substrate, with the diamond facing
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Figure 13 Calculated transmission for a diamond film on glass for rms surface roughness values of 10, 20, and 40 nm.
the chalcogenide. The sandwich is then warm pressed, during which most of the chalcogenide is squeezed out, leaving a thin film that bonds the rough diamond surface to the ZnSe. Finally, the silicon is etched away, leaving the smooth side of the diamond film exposed and the rough side of the diamond film immersed in the bonding glass. As another example, Fuentes et al. [90] have described a method to bond diamond films anodically to glass substrates. Diamond was found not to bond anodically directly to glass rings, ˚ thick, was sputtered on the surface so an interfacial layer of SiC, 1000 A of diamond films, which had been previously deposited on silicon substrates. The Si-diamond substrate was then placed with the diamond against glass, and a combination of heat and electrical bias was used to accomplish the bonding. Subsequently, the silicon was removed by etching, leaving diamond membranes attached to glass rings. Also, a method for attaching diamond to infrared-transmitting glass using a commercial optical glue has been reported by Ran et al. [91]. The majority of studies of diamond-coated optical substrates, however, have involved deposition of the diamond onto the substrate in a CVD growth chamber. The rest of this section is concerned with examples that use this direct growth approach. CVD methods have been used to deposit diamond films on a variety of substrates of optical interest including quartz [92,93], fused silica [94–
Diamond Film Optics
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96] sapphire [93,96,97], soda-lime glass and lead glass [98], borosilicate glasses [99–101], germanium [96], MgO [98], and ZnS [97]. Silicon, one of the most common substrates used for CVD diamond deposition, is also of interest for certain optical applications including detectors. For some optical substrates, CVD of diamond is problematic because of (1) low temperature requirements, (2) coefficient of thermal expansion (CTE) mismatches with diamond, or (3) the substrate being damaged by the deposition plasma. All three of these concerns are present, for example, with ZnS and ZnSe. These materials would be severely etched by the atomic hydrogen present in typical diamond deposition plasmas. Also, chemical considerations require that the substrate temperature during deposition be less than 650⬚C. Finally, the CTE of these substrates is several times higher than that of diamond. In spite of these concerns, successful diamond deposition on these substrates can be accomplished by first applying a protective interlayer, as described by Costello et al. [97] which protects the substrate from the deposition plasma [102]. In contrast, silicon exhibits none of these concerns to a serious degree. A silicon substrate is compatible with direct contact with deposition plasmas, temperatures over 1000⬚C can be tolerated, and although the CTE of silicon is not identical to that of diamond, the differences are such that good adhesion can be obtained over a wide range of deposition conditions. Most optical substrates fall somewhere between the extremes of ZnS and Si, with one or more of the preceding three concerns being of issue. An optical substrate of considerable interest for diamond coating is silicate-based glass. Such glass, arguably the most common optical material, lacks the resistance to erosion and abrasion exhibited by diamond. Consequently, direct CVD of diamond on common commercial glass is of interest for enhancing resilience to hostile environments. Issues associated with such deposition depend on the glass type. Table 3 shows the composition of several common glasses, including soda-lime, borosilicate, and lead glass,
Table 3 Comparison of Commercial Silicate Glasses Showing Composition by Weight %, Strain Temperature Ts, and Room-Temperature Coefficient of Thermal Expansion Glass type
SiO2
B2O3
Al2O3
CaO
MgO
Fused silica Soda lime Borosilicate Alkali lead
99.9 72.6 81.0 77.0
0.8 13.0
1.7 2.0
4.6
3.6
Source: Refs. 103 and 104.
1.0
PbO
Na2O
8.0
15.2 4.0 9.0
K2O
Ts (⬚C)
CTE (10⫺7/⬚C)
5.0
956 473 510 390
5.5 92 33 93
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as well as the CTE in each case* [103,104]. Soda-lime glass is commonly used for windows, containers, and lamps; borosilicate glass is used for cookware, labware, and headlamps; and alkali lead glass is used for lamp tubing and sealing applications. Of principal interest is the CTE, which ranges from a value somewhat below that of diamond for the case of fused silica to values an order of magnitude higher than diamond for the case of soda-lime and lead glasses. Also varying greatly with glass composition is the effect of temperature on the mechanical properties, with the glass changing gradually from an elastic solid at room temperature to a viscous liquid at high temperatures. A temperature of relevance to deposition applications is the strain temperature, which corresponds to a viscosity of 1014.5 poise. Generally speaking, internal strain cannot be introduced into the glass at temperatures below this value. In order to minimize internal strain in the diamond film on the cool-down after growth, the deposition temperature should be less than the strain temperature of the glass. Considering the compositions in Table 3, the strain temperatures vary from 956⬚C for fused silica to less than 500⬚C for the soda-lime and lead glasses. In any case, for silicate-based glasses the principal matrix seen by the diamond film is that of SiO2. The interface between CVD diamond films and SiO2 and between CVD diamond films and Si has been studied by Pickrell et al. [95] by analyzing the interface side of flakes of diamond film removed from the substrate. In the case of conventional silicon substrates, several groups have reported that an essentially continuous SiC layer forms during the initial growth [95,105,106]. However, in the case of SiO2 substrates, silicon carbide was observed only in discrete, particulate-like areas. Both Si and SiO2 will react with C to form SiC at elevated temperatures. However, in the reaction of C with SiO2 a gas, CO, is evolved. The silicon carbide can also react with the SiO2 substrate to form more gases, as [95] 2SiO2 ⫹ SiC = 3SiO ⫹ CO
(24)
where both products are gases. The extent of reaction between arriving carbon and silica to form silicon carbide will be limited by the diffusion of gaseous species away from the interface, a process that will slow down as the diamond film grows, eventually reaching a local equilibrium. As suggested by Pickrell et al. [95], the crystals of silicon carbide can act as nucleation sites for diamond. However, gas evolution may impede the growth of diamond parallel to the substrate, which produces a noncontinuous, par-
*In terms of commercial codes, the glasses in Table 3 correspond to the following Corning numbers: fused silica (Corning 7940), soda lime (Corning 0080), borosilicate (Corning 7740, also called ‘‘娃Pyrex’’), alkali lead (Corning 0010).
Diamond Film Optics
643
ticulate type of intermediate SiC layer at the interface. Other potential nucleation sites on silica surfaces include residual diamond crystals from seeding procedures, surface scratches, and possibly a ternary Si-O-C phase [107]. In spite of the evidence for a patchy SiC interface, block-on-ring tribotests have indicated that diamond films adhere well to fused silica substrates [94]. Diamond deposition on fused silica can be accomplished using the normal range of CVD diamond deposition temperatures, typically between approximately 700 and 900⬚C. However, for the more common glasses, it has been necessary to develop lower temperature deposition systems. In the case of microwave plasma-assisted CVD, for example, lower substrate temperatures have been accomplished by a variety of methods including [92,94,99,100,108–118]: decreasing the microwave input power and lowering the pressure such that the size of the plasma increases, thereby decreasing the plasma power density; cooled substrates; pulsed power systems; and magnetoactive plasmas, including electron-cyclotron-resonance (ECR) plasmas. Although most low-temperature deposition reports have involved microwave plasma CVD systems, other deposition systems have been adapted to low-temperature deposition, including hot-filament systems [119], combination ECR/filament systems [120], and rf plasma systems [121]. Input gas mixtures that have been optimized for high-temperature deposition are not generally optimum for low-temperature deposition. In many, but not all, of these cited cases, oxygen-containing gases were added to hydrogen-methane mixtures, including CO, CO2, and O2. In some cases, H2 has been eliminated from the input gases [114], and in other cases alternate sources of carbon, including fullerenes, have been investigated [122]. A hindrance to low-temperature deposition is the fact that abstraction of hydrogen from growth sites is a temperature-activated process, generally involving hydrogen atoms. As a result, growth rates at low temperatures exhibit an experimental activation energy of approximately 0.6 eV [99], consistent with a model by Harris and Weiner [123] in which radical surface site pairs are involved in the incorporation of new carbon species on the surface. Consequently, low-temperature (less than 500⬚C) CVD diamond deposition that results in uniform coverage of areas of significant size show rather low growth rates, on the order of 0.1 m/hr. Also of interest as feed gases for low-temperature deposition are halogens [124]. The dissociation of molecular fluorine into fluorine radicals occurs at rather low temperatures and fluorine may be beneficial in abstracting hydrogen from the growth surface at low temperatures. As would be expected from the discussion in Secs. III.A and V.A, high transparency of diamond-coated substrates requires a smooth diamond surface. Two examples of diamond-coated borosilicate glass are shown in Fig. 14. Both samples in this figure consist of polycrystalline CVD diamond
644
Reinhard
Figure 14 Photograph of two diamond-coated borosilicate glass samples in the as-grown state. The sample on the left is a large-grain film with appreciable scatter and low transmission in the visible. The sample on the right is a highly transparent, small-grain diamond film [100].
deposited on 5-cm-diameter glass substrates by microwave plasma-assisted deposition and both samples are in the as-grown, nonpolished state [99]. The difference in visible transmission between the two samples is striking. Light transmission through the sample on the right-hand side of the photograph allows good visibility of the optical bench. Holes in the bench that are approximately a meter distant from the sample are clearly noted. In contrast, the sample on the left has sufficient scattering that the holes in the optical bench cannot be seen. Both samples are of comparable thickness—1.4 m for the sample on the right and 1.7 m for the sample on the left. However, they differ in terms of the initial seeding method and therefore the final grain size and surface roughness. The sample on the left is representative of films with grain sizes on the order of a micrometer and with surface roughness rms values on the order of 100 nm. As a result of the surface roughness, there is strong scattering of visible light and the sample displays a frosted glass appearance. Such samples appear translucent rather than transparent;
Diamond Film Optics
645
objects are clearly visible only if they are directly adjacent to the diamond film. The sample on the right is representative of finer grain films where the average grain size is substantially less than a micrometer and the surface roughness is measured in tens of nanometers. For the particular sample illustrated, the surface roughness was experimentally measured to be in the range of 20 to 30 nm. Although the grain size is small, the Raman spectrum, shown in Fig. 15, shows the characteristic diamond peak at 1332 cm⫺1 and does not indicate appreciable nondiamond carbon content. Figure 16 shows the measured transmission through the two samples shown in Fig. 14, with a 1-cm-diameter detector located 2 cm behind the sample and a 0.4-cm-diameter beam incident on the sample. For the largegrain film, the measured transmission is less than 10% throughout the visible. For the small-grain film, the transmission falls from about 70% in the red portion of the visible spectrum to approximately 40% in the blue. Comparing these results with the calculated transmissions in Fig. 14, the fall-off in experimental transmission with decreasing wavelength is consistent with a surface roughness value of approximately 35 nm, slightly larger than the surface roughness as measured by a stylus. The calculated transmissions of Fig. 14 predict that a surface roughness of 10 nm or less should result in high transmission throughout the entire visible range. Such transmission is, in fact, achievable for as-grown polycrystalline CVD diamond films on glass, as shown in Fig. 17 for a 0.43-m-thick diamond film deposited on Corning 7059 borosilicate glass [125].*
C.
Stress and Adhesion Considerations
Film adhesion can be an important issue for CVD diamond films on optical substrates because of stress in the films. This is particularly true for substrates whose thermal expansion coefficients are poorly matched to that of diamond. For example, Raman analysis of diamond films on soda-lime glass have shown a shift in the diamond peak from 1332 to 1340 cm⫺1 [98] with the shift to higher wavenumber indicating that compressive film stress is present. Large stresses can cause sample bowing and facilitate film delamination from the substrate. In fact, the phenomenon has been proposed as a means of forming freestanding diamond films, based on a study in which diamond deposited by hot-filament CVD onto Corning 7059 glass was observed to separate spontaneously from the substrate [101]. For most appli-
*Corning 7059 is a low sodium content borosilicate glass, otherwise similar to 娃Pyrex. It is used in electronic applications, including flat panel displays.
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Figure 15 Raman spectrum for the transparent diamond film shown on the righthand side of Fig. 14.
cations, however, it is desired that the film remain intact to the optical substrate. The dependence of stress in diamond films has been studied by Windischmann et al. [126] as a function of gas composition and deposition temperature. If diamond is deposited on a substrate whose CTE is higher than that of diamond, a compressive contribution to the total stress in the film is developed as the film-substrate combination is cooled from the deposition temperature to room temperature. In addition to the thermally induced stress caused during cool-down, diamond films have an intrinsic tensile stress that depends on methane concentration, among other things. Its origin is explained in terms of a grain boundary relaxation model due to a fine-grained porous microstructure [126]. Consequently, diamond films may be under either tensile stress or compressive stress, depending on the substrate and on deposition conditions. If thermally induced compressive stress and intrinsic tensile stresses cancel, the resulting films have very low net stress. For numerical calculation purposes, the thermal compressive stress may be estimated by breaking the temperature into i intervals and calculating a stress for each interval according to [126]
i =
冉 冊 E 1⫺
(␣F ⫺ ␣S)⌬Ti
(25)
where E/(1 ⫺ ) is the biaxial modulus and (␣F ⫺ ␣S) is the average difference in CTE values for the film, ␣F, and the substrate, ␣S, at the midpoint of the ith temperature interval, ⌬Ti. The total thermal stress, th, is calculated by summing the i. An estimate of the internal tensile stress, in, is given by Windischmann et al. as
Diamond Film Optics
647
Figure 16 Measured transmission in the visible and near infrared for the two samples shown in 14. (a) The large-grain film and (b) the small-grain film [125].
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Figure 17 Measured optical transmission for a 0.43-m-thick diamond film on borosilicate glass (Corning 7059) that shows high transmission throughout the visible and near infrared [125].
in =
冉 冊冉 冊 E 1⫺
␦ d
(26)
where d is the grain size of the polycrystalline film and ␦ is estimated to be 0.077 nm for diamond [126]. The total stress, tot, is equal to the difference between th and in and either compressive or tensile, depending on the sign. For silicon, the thermal stress is always compressive. Although the CTE of silicon drops below that of diamond at high temperatures, the difference is not enough to counteract the larger CTE of silicon at low temperatures. For diamond films on silicon with grain sizes of the order of 0.1 m and larger the combined stress is compressive, based on both calculations and experiments that measured the bowing of diamond diaphragms produced by back etching the silicon. In this regime, as the grain size increases, other things being equal, the net strain would be expected to increase. This is consistent with a report by Gamlen et al. [127], who investigated diamond film adhesion on silicon by measuring delamination after Vickers indentation and found the delamination diameter to increase with increasing grain size. For very fine grain diamond films on silicon, however, the net stress is generally tensile.
Diamond Film Optics
649
Figure 18 shows the CTE values for diamond and Corning 7059 glass as a function of temperature. The glass CTE value is fairly constant until near the strain point of 593⬚C, at which it increases sharply. Consider, as an example, a diamond film on 7059 glass cooling down from 475 to 25⬚C. Taking the biaxial modulus for diamond to be 1345 GPa [128], the thermally induced strain may be estimated from Eq. (25) to be approximately 1.7 GPa compressive, using a ⌬Ti value of 100⬚C. Consequently, from Eq. (26), films with grain sizes larger than approximately 70 nm would be in a state of compressive stress and films with smaller grain sizes would be in a state of tensile stress. Grain size depends most directly on seeding density; however, for a given seeding density, grain size generally increases with increasing film thickness. Equations (25) and (26) are useful for determining trends. Neglected, however, are effects such as structural relaxation in the substrate [129] and compressive stress due to residual hydrogen and sp2 carbon [126]. In experimental studies of diamond on glass, some films have shown good longterm adhesion with, for example, diamond films on both Corning 7059 and 7040 substrates that pass the tape test for adhesion several years after deposition and after repeated thermal cycling to 200⬚C. Other films, however, have shown spontaneous delamination after times ranging from several minutes to several days after removal from the deposition chamber. Adhesion
Figure 18 The coefficient of thermal expansion coefficients for diamond and Corning 7059 glass.
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Table 4 Film Thickness Dependence for Diamond Film Adhesion to Glassa Thickness (m)
Adhered?
0.7 1.0 1.7 2.5
Yes Yes Yes No
a
Substrates are 5-cm-diameter Corning 7040 (娃Pyrex) glass.
trends for translucent diamond films on borosilicate glass such as pictured on the left side of Fig. 14 are shown in Tables 4 and 5 with adhesion becoming more problematic as both substrate temperatures increase and as film thickness increases. These trends are consistent with general expectations from Eqs. (25) and (26), given the CTE values of the glass substrates. Long-term adhesion of such films on 7059 glass was reported when diamond film thicknesses were less than 3 m and deposition temperatures were less than 480⬚C [100]. VI.
PHOTON DETECTION AND EMISSION
This section briefly considers diamond’s role in electro-optics. Both radiation-induced electrical conductivity within diamond and photon emission from diamond have been used extensively for diagnostics of diamond electronic states. The diagnostic use of such phenomena is not within the scope
Table 5 Substrate Temperature Dependence and Thickness Dependence of Diamond Adhesion to Glassa Pressure (torr)
Temperature (⬚C)
Thickness (m)
Adhered?
15 13 11 9
505–510 480–483 445–455 440
2.3 2.1 1.3 1.0
No No Yes Yes
a
Substrates are 5-cm diameter Corning 7059 glass. The substrate temperature was measured by optical pyrometry.
Diamond Film Optics
651
of this chapter but is treated in Chap. 3. The focus here is on applicationrelated issues. To the extent that photon absorption in diamond involves electronic transitions, the electric properties of diamond are altered. Photons with energies greater than the band gap, that is,with energies greater than 5.47 eV, will with some probability generate free carriers in diamond that cause a photoconductive effect. Electronic transitions involving states in the gap can also contribute to photoconductivity in diamond for photons of lessor energy. Likewise, electronic transitions between states, whether band to band or involving gap states, can with some probability involve photon emission. Consequently, both photon detection and photon emission applications are of interest. The use of single-crystal diamond for photon detection has a long but not particularly extensive history with early work on diamond photoconductivity being reported in 1923 [130]. However, as is the case for many other applications, the advent of CVD diamond provided an impetus for renewed interest. The wide band gap of diamond makes it of particular interest for UV detection [131–135]. An ultraviolet detector should preferably have no sensitivity to visible light, a large sensitivity to UV light, and a small dark current. A useful figure of merit is the quantum efficiency, QE, taken as the total number of electrons collected at the contacts divided by the number of incident photons, or as QE =
Iphh Poptq
(27)
where Iph is the photo current, h is the photon energy, and Popt is the optical power. The simplest UV detector structure is a two-terminal device consisting of diamond between two metal electrodes, an example of the so-called metal-semiconductor-metal (MSM) structure. The quantum efficiency for an MSM device is related to diamond material properties and the electrical bias according to [136] QE =
F () L
(28)
where F is the electric field applied between the metal contacts, L is the separation between the contacts, is the carrier mobility, is the carrier lifetime, and is the fraction of incident power absorbed in the diamond,
= (1 ⫺ R)(1 ⫺ e⫺␣t)
(29)
In an early study of CVD diamond MSM UV detectors by Binari et al. [133], a comparison was made between the performance of a natural diamond type IIa detector and that of a CVD diamond detector. In both
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cases, the MSM structure consisted of two evaporated aluminum stripes separated by 350 m on the surface of the diamond. Both devices showed an acceptable dark current, approximately 10⫺12 A, and both showed a substantially larger UV response than visible response. The quantum efficiency of the type IIa detector at a UV wavelength of 200 nm was approximately three orders of magnitude higher than the quantum efficiency in the blue portion of the visible spectrum (400 nm) and the quantum efficiency of the CVD diamond detector was approximately two orders of magnitude higher at 200 nm than at 400 nm. However, the two detectors differed significantly in the UV 200-nm quantum efficiency, which, when biased at 100 V, was 39% for the type IIa diamond and 0.09% for the CVD diamond. As is evident from Eq. (28), the quantum efficiency of an MSM photodetector is strongly dependent on the product of the sensor material. Polycrystalline materials in general tend to exhibit lower mobilities and lower lifetimes than their single-crystal counterparts and this is also the case for diamond. Therefore, a materials challenge for CVD diamond photodetectors is the achievement of higher products [137]. For a given bias voltage, the QE will also improve with decreasing electrode spacing. Using MSM structures with 20-m inter-electrode spacing, L, Chan et al. [134] have fabricated devices with five orders of magnitude difference in sensitivity between deep UV light and visible light. Another materials-related issue for CVD diamond UV detectors is persistent photoconductivity, which results in a memory effect that can have significant time constants, on the order of 1000 sec or longer. The origin of this effect is that electrons excited by the UV irradiation are trapped in deep traps that have long release times. Detectors can be reset, or bleached, by heating them sufficiently to release the trapped carriers. This phenomenon can be put to use, as proposed by Gonon et al. [138] for dosimetry, by measuring thermally simulated currents after high-energy photon absorption. However, for typical UV detection applications a persistent photoconductivity is undesirable. Chan et al. [134] reported a postgrowth methane-air treatment that reduces the duration of the persistent photoconductivity from in excess of 500 sec to less than 5 sec. Diamond electronic device structures other than the MSM, for example, Schottky barrier diamond photodiodes, are also applicable for photon detection [139,140]. A related application is photoconductive switching. Single-crystal type IIb diamond has been evaluated as a detector for several pulsed laser sources by Young et al. [141] and shown to have subnanosecond response times. Subnanosecond, kilovolt switching of microjoule optical pulses has been reported in insulating single-crystal diamond [142]. In the range of 0.22– 0.6 m, defect and impurity photoabsorption provides the switching mechanism. As noted by Angus et al. [131], CVD diamond may have significant
Diamond Film Optics
Figure 19 Ref. 143.)
653
Typical waveform of a CVD diamond photoconductive switch. (From
advantages over natural diamond because of reproducible levels of vacancy and nitrogen defect complexes as well as increased simplicity of doping. Using CVD diamond, Aikawa et al. [143] demonstrated subnanosecond detection of light pulses from a KrF laser as shown in Fig. 19. With regard to electro-optic applications that involve photon emission, band-to-band recombination events in diamond would not be efficient sources of photons because diamond has an indirect band gap. As discussed by Yoder [144], however, another approach may be considered for diamond in that impurities may be intentionally added in order to add states within the large forbidden energy gap so as to provide optical transition energies. In fact, visible-light lasting has been demonstrated in single-crystal diamond by external pumping of electrons into such color centers [145]. Also, diamond based p-n diodes and Schottky diodes have demonstrated electroluminescence (EL) upon forward bias [146–148]. In these cases, the EL and cathodoluminescence spectrum coincide and reflect the impurity and defect content of the diamond. Fahrner et al. [149] used nitrogen implantation into natural type IIa single-crystal diamond and showed a modification of the EL spectrum.
VII.
SUMMARY
The unique intrinsic optical properties of diamond are realizable in polycrystalline CVD diamond in both thick- and thin-film form, so the material is of interest for optical applications involving photon transmission, photon
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detection, and photon emission. With regard to transmission, special care is needed to avoid scattering caused by rough surfaces. Smooth surfaces are achieved either by polishing, by fine grains, or by oriented growth. Also necessary for high transmission are growth conditions that contribute neither appreciable impurities nor appreciable nondiamond carbon [150]. For thick films containing large grains, the mechanical strength is significantly less than for single-crystal diamond because of the large flaw sizes inherent in such structures. However, the strength is still greater than that of many other contending materials. The attractive surface attributes of diamond, including hardness and chemical resistance, may be added to more common optical substrates by thin-film diamond coatings [151]. For many substrates, stress in such films is controllable by balancing compressive thermal stress and tensile internal stress. As an indirect band gap material, photon emission heavily involves impurity and defect states. Such impurities can be intentionally added to provide color centers for the enabling of optical transitions. The large band gap allows ultraviolet photon sensing with essentially visibleblind response. Viability of CVD diamond for optical applications is tied to (1) deposition technology and the ability to achieve simultaneously useful growth rates and the necessary optical properties, (2) surface roughness values, (3) mechanical strength considerations, and (4) compatibility between CVD diamond and optical substrates.
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Index
Ablation, laser, 333–335, 354–358 Abrasive wear, 439, 440 Absorption, optical band-to-band, 609 defect induced, 624–626 due to C — H bonds, 622, 632 impurity related, 613–616, 621–624 infrared, 612, 622–624, 632 intrinsic, 613 multiphoton, 336 resonant, 376 ultraviolet, 609 Urbach tail, 611 weak tail, 611 Accademio del Cimento, 18 Acceptors (see Boron) Acetylene, 66, 67, 109, 164, 165, 214, 247, 269–272, 288, 296, 304 Acoustic applications, 12 Actinometry, 93, 273 Activation energy, 250 Adhesion diamond to substrate, 437, 458–475, 645–650 metal to diamond, 556 treatments, 476
AES (see Auger electron spectroscopy) AFM, 551 Allemanna Svenska Elektrisia Aktiebolaget, 18, 20 Alloys high-silicon aluminum, 425, 444, 475, 482, 484 low-silicon aluminum, 478, 484 ␣ parameter, 38, 81, 207, 252–260, 264, 265, 266 Alumina, 519 Aluminum (see also Alloys) as electric contact, 586, 599 as via fill, 556 Aluminum nitride, 568 Amorphic diamond, 354, 355 Amorphous carbon, 342, 354, 462 Angstrom’s method, 536 Anharmonic pumping, 380 Annealing, 590, 59 Antireflection diamond, as, 634 layers, on diamond, 636 Applications of diamond acoustic, 12 cold cathode emitter, 12 665
666 [Applications of diamond] cutting tools, 10, 137, 141 (see also Tools, cutting) detectors, 12, 651, 652 diodes, 13, 580, 652 electro chemical, 13 electrodes, 10 electron and photon sources, 11 electronic active devices, 10, 579–602 passivation layers, 10 microelectronics, 14 packaging, 10, 561–566 gyrotron windows, 10, 11, 627 heat sinks, 10 lasers, 10, 11, 607, 627 nuclear particle dectors, 10 optical coatings, 10, 11, 637–648 optical windows, 10, 11, 627 sensors, 10, 12 speakers, 10 surface acoustic wave, 10, 12 thermal management, 11, 137, 141, 513–522 thermistors, 1, 10 transistors, 13, 579–602 transmission (see also Transmission, optical), 11 ultraviolet dectors, 10, 12, 651, 652 vacuum microelectronics, 12 wear surfaces, 10, 439–441 x-ray masks, 11 Applied Science and Technology (see ASTek, Inc.) Applied Sciences Inc., 532 Arcjet CVD (see also D-c arcjet CVD), 5, 56, 59, 66, 93, 72, 141–210, 527 Arrhenius, 249 Arsenic as dopant, 584 ASTeX Inc., 229 Atomic density of diamond, 5 Atomic Force Microscopy (see AFM) AT&T Bell Laboratories, 545
Index Auger spectroscopy, 35, 78, 343, 364, 621 Bachmann diagram, 242, 243, 248 Balagia’s figure of merit, 579, 580 Band gap (see also Energy gap) direct gap, 611 indirect gap, 610 Bell Laboratories, 21 Beryllium oxide, 519 Bias enhanced nucleation, 237, 238, 524 Binder, metallic, 426, 460, 470, 474 Bipolar transistor, 585 Blackbody radiation subtraction of, 413 Boltzmann’s constant, 374 Boltzmann distribution, 380 Boron activation energy of, 582 as acceptor, 580, 586, 589–592, concentrations of, 590 implantation of, 592 in Hope diamond, 615 in type IIb diamond, 614 ionization energy, 580 optical effects of, 613, 621 Boudouard reaction, 389 Boundary layer, 156–159 Brazing, 430, 450 Breakdown, electric field, 579 Bulk modulus, of diamond, 520 12
C (see Carbon-12) C (see Carbon-13) C — H stretch modes, 622 C2, 96, 97, 160, 161, 162, 269, 389 C2H2 (see Acetylene) Carat volume, 5 weight, 5 Carbon, non-diamond amorphous, 454, 525 atomic, 106, 108, 143, 164 carbon black, 353, 362 fullerene, 525 13
Index [Carbon, non-diamond] glassy, 352 graphitic, 525 phase diagram, 2, 19, 338, 350 Carbon-12, 352, 364, 412 Carbon-13, 352, 364, 412 Carbon conversion efficiency, 142, 184– 187, 223, 225, 250, 251 Carbon tool steels, 426 Carbonaceous precursors, 341, 364 Carrier concentration, 581 mobility, 9, 581 CARS (see Coherent anti-Stokes Raman spectroscopy) Case Western Reserve University, 22, 23 Cast iron milling, with diamond, 497 polishing diamond with, 630 Catalysts, 350, 525, 532 Cathodeluminescence, 34, 653 Cemented carbides, 426 Ceramic bisque, 442 green, 442, 447 CH3 (see Methyl) CH4 (see Methane) Chemical resistance, of diamond, 439, 520, 530 Chemical vapor deposition (see also Arcjet, Combustion, Laser, Hot filament, Microwave), 19, 22 high-pressure, 18 laser excited, 5, 378 low-pressure, 73 thermal plasma, 141 Chlorine, 267, 268 C/H/O diagram, 242, 243 Circuits, diamond, 594–598 Cleavage strength, of diamond, 7 Cleavage, through grains, 457 CO as carbon source, 270, 288, 364, 365 energy levels of molecules, 378 laser spectrum, 387
667 [CO] laser transitions, 378, 379 optical pumping of, 386 single crystal, 410 vibrational mode, 388 Cobalt, 426, 460, 463 as crystalline substrate, 343 Coherence length, 331 Coherence time, 330 Coherent anti-Stokes Raman spectroscopy, 89, 91, 95, 132, 133, 275, 276, 278, 288 Coherent light, 328–332 Color, of diamond blue, from boron, 613 brown, from nitrogen, 615 brown, from sp2 bonds, 626 green, from irradiation, 616 pink, from manganese, 616 yellow, from nitrogen, 613–615 Color centers, 653 Combustion CVD, 5, 56, 68, 96–98, 109, 303–323 Competitive growth, 76 Composites carbon-epoxy, 488 graphite-copper, 534 metal-matrix, 425, 444, 479, 532, 533 Compressibility of diamond, 6 Conduction band, 610 Conductivity electrical of diamond (see also Electronics, diamond), 7, 49 thermal (see Thermal conductivity) Contacts (see also Schottky) ohmic, 584, 591, 599 Contamination by filament, 621 in feedgas, 621 Conversion factors, energy, 361 Copper as crystalline substrate, 343 as electrical contact, 600 as heat sink, 567 as via fill, 556 Correlation function, 329
668 Cost, of diamond, 341 Crack propagation, 464 velocity, 464 CRDS (see Ring-down spectroscopy) Crystalline Materials Corp, 555 CTE (coefficient of thermal expansion), 529 of diamond, 7, 436, 437, 518, 530, 649 of glass, 641, 649 Cubic boron nitride, 429 Cubo-octohedral diamond, 368 Cullinan diamond, 615 Cutting, of diamond, 450, 540, 541 Cutting tools (see Tools) CVD (see Chemical vapor deposition) DARPA, 561 D-c arcjet anode, 149, 152 cathode, 149, 152, 153 CVD, 59, 68, 70, 92, 98, 106 instabilities, 158 nozzle, 154–156, 160 DeBeers, 449 Defects, in diamond absorption spectra effects, 624–626 interfacial, 464 Degree of mutual coherence, 329 Delamination, 460, 468, 645, 648–650 Density, of diamond, 5, 520, 530 Detectors ultraviolet, 651, 652 Diamond classification Type Ia, 613–615 Type Ib, 613–615 Type IIa, 592, 613–615 Type IIb, 613–615 Diamond grit, 365 Diamond like carbon, 338, 342, 486 Diamond properties band-gap, 13, 610, 611 boron doped (see also Boron), 13 cleavage strength, 7 compressibility, 6
Index [Diamond properties] density, 5, 520, 530 electrical conductivity, 7, 49, 581 energy gap, 39, 579, 610–613 Hall mobility, 8 hardness, 6, 433, 434, 530 optical transmission (see also Transmission, optical), 8, 9 refractive index, 8, 609 resistivity, 7, 514, 519, 520, 530 specific heat, 7, 520 thermal conductivity (see also Thermal conductivity), 7, 9, 13, 43, 45, 46, 49 thermal diffusivity, 45 Diamond seeds, 352 Diamond synthesis (see Hot-filament CVD, Arcjet CVD, Combustion CVD, Microwave CVD, rf plasma CVD, Laser-assisted CVD) Diamonex, 137 Diaphragms, diamond, 627–629, 631– 636 Diatomic molecule energy storage, 326, 376 single-quantum processes, 382 Dielectric, diamond breakdown (see Electrical breakdown) constant, 520, 530, 579 loss tangent, 520, 530 strength, 530 Diffusion of boron, 590 of carbon, 345 Dipole moment, 382 DLC (see Diamond like carbon) Domes, diamond, 608 Doner, 584 Dopant Electronics Materials Research, 555 Doping, of diamond, 586, 590, 598–600 Drilling, of diamond, 540, 541, 555, 556 Drilling tools (see Tools, rotary) Dymalloy, 531–531
Index EEDF (see Electron energy distribution function) EELS (see Electron loss spectroscopy) Effective mass, 580 Electric breakdown field, 530 Electro-chemical, 13 Electroluminescence, 653 Electron energy distribution function (EEDF), 256–288 Electron density, 278, 293, 294 Electronics, diamond, 10, 141, 579– 602 Electron loss spectroscopy (see Film characterization) Electron-photon interactions, 344 Electro-optics (see Optoelectronics) Emission spectrum, 399 End mills (see Tools) Energy, conversion factors, 361 Energy cost (see Joules per carat) Energy dispersion spectrum, 392, 395, 402 Energy distribution Boltzmann, 380 Treanor-Rich-Rehm, 380 Energy gap, 39, 579, 610–613 Energy transfer (see also Vibrations) Epitaxial heteroepitaxial, 79, 80, 461 homoepitaxial, 1, 75, 148, 581, 589, 594 Erosion rain resistance, 634, 635 sand resistance, 635 Etching, of diamond ECR plasma, 631 ion beam, 584 mesa, 592 reactive ion, 550, 594 Ethylene, 339, 362 Excimer (see Lasers) Excitation electronic, 326 rotation, 326 vibration, 326
669 Failure modes of tools, 448 Ferrous, machining by diamond, 497 FET (see Field effect transistor) Field effect transistor, 582–599 Filament carburization, 124 contamination, 125 materials rhenium, 122, 128 tantalum, 122, 128 tungsten, 58, 119, 122, 128 poisoning, 101, 102, 129 sooting, 129 Film characterization, 21–54 Auger electron spectroscopy, 35, 78, 343, 364, 621 cathodoluminescence, 34 electron loss spectroscopy, 30, 33, 34, 78 high-resolution electronmicroscopy, 33 laser microscope mass analysis spectroscopy, 35 low-energy electron diffraction (LEED), 21, 35 Raman spectroscopy (see also Raman), 28, 148, 150 Rutherford backscattering, 36, 78, 127 scanning electron microscopy (see also SEM), 28, 30, 40, 78, 180, 182, 184, 201 scanning force microscopy, 47–49 scanning probe measurements, 47, 48 scanning thermal microscope, 49, 50 scanning tunneling microscopy, 47, 48 secondary ion mass spectrometry, 35, 36, 37 thermal conductivity measurements (see also Thermal conductivity), 46 conventional method, 47 thermo-optical method, 47 transmission electron microscopy, 28, 30, 32, 33, 78, 340, 342
670 [Film characterization] X-ray diffraction, 28, 36, 37, 40, 41, 78, 264 Finite element analysis, 288, 558 Flame temperature, 98 Flip-chip, 557–560 Flourine, 267–268, 526 Fourier analysis, 332 Fourier transform infrared analysis, 95, 98, 132, 269 Fracture, modes of, 457 Fracture toughness, of diamond, 435, 438–441 Freestanding diamond, 177, 627–629 Friction, coefficient for diamond, 433, 434, 499 FTIR (see Fourier transform infrared analysis) Fullerene, 525 Gallium arsenide, 569 Gas chromatography, 88 Gem diamond, 609 General Electric, 2, 17, 18, 20, 326, 352 Germanium, as substrate, 354, 355 Glass (see also Substrates) alkali lead, 641, 642 borosilicate, 641–650 TM Pyrex, 650 soda lime, 641, 642 Gold as electric contact, 590, 591, 599 Grain boundaries, in diamond, 458, 553 Grains cleavage through, 457 in PCD, 427 Graphite disordered, 340 machining, 486 microcrystalline, 343 transformation of diamond to, 636 Growth, of diamond on diamond seed, 405–409
Index Growth rate, 71, 72, 141, 142, 147, 148, 173, 174, 176, 181, 185, 186, 190, 201, 207, 211, 225, 242–247, 249–252, 260–265, 341 calculation of, 373, 374 Gyrotron (see Windows, diamond) Hall mobility, of diamond, 8 Hard coatings, 440 Hardness, of diamond, 6, 433, 434, 530 Harvard University, 18 Heat capacity of diamond, 7, 516 Heat exchangers, 533 Heat sink, 10, 512, 516, 518, 529, 540, 565 Heat spreader, 513–522 Hewlett Packard, 561 High-pressure, high temperature diamond, 3, 18, 326, 349, 352, 586 High-resolution electron microscopy, 33 Hole concentration, 580 mobility, 581 Homoepitaxy, 1, 67, 75, 79, 80, 148, 266 Hope diamond, 615 Hot-filament CVD, 5, 24, 56, 58, 66, 67, 68, 70, 75, 88, 89, 98, 99–105, 119– 140, 368, 527, 528 HREM (see High-resolution electron microscopy) Hydrogen abstraction, 64, 65, 67, 69, 238, 239, 240, 268 atomic, 59, 60, 64, 65, 71, 72, 87, 89, 93, 95, 100, 101, 104, 105, 108, 142, 143, 146, 160, 161, 170, 213, 214, 238, 239, 240, 242, 247, 262, 273, 274, 276, 356, 363
Index [Hydrogen] dissociation, 58, 102, 106, 121, 128, 143, 149, 157, 213, 238, 291, 292 in polycrystalline diamond, 622 molecular flow rate, 370 optical effects, 621 p-type diamond, 594, 599 recombination, 105, 125 surface passivation, 369 IBM, 558 Impurities absorption spectra effects, 621–624 in microwave CVD diamond, 622 Incoherent light, 328–332 Infrared absorption (see Optical absorption) Infrared diode laser absorption spectroscopy, 88 Infrared laser absorption, 130 Infrared spectrum gas phase, 396 Insulated gate, 588 Interface (see also Adhesion, Microstructure) chemical bonding, 462 degradation, 476 liquid-solid, 352 Interference fringes, 331 Ion implantation, 343, 585, 591, 592 Ionization energy of boron acceptors, 580 of molecules, 359 Iron, as catalyst, 350 Johnson’s figure of merit, 579, 580 Joules per carat, 341, 350, 376, 408, 416 Kennametal, Inc., 506 Key’s figure of merit, 579 Knoop hardness of several optical materials, 360 Knoop indenter, 6
671 Langmuir probe, 279 Laser chemical vapor deposition, 366– 369 Laser driven reactions, 327, 370 Laser excited chemical vapor deposition (see also Optical pumping) CO lasing lines used in, 392 experimental setup, 391 Laser-induced fluorescence, 89, 91, 96, 97, 110, 132, 160, 161, 273, 275, 276, 278 Laser microscope mass analysis spectroscopy, 35 Laser physical vapor deposition (see Ablation) Laser-plasma technique, 341 Laser pumping, 336, 337 Laser pyrolysis, 333 Lasers argon ion, 352 CO (see also CO), 379 CO2, 348 excimer, 348, 371, 547 for diamond synthesis, 325–417 high power, 607, 627 Nd-YAG, 342, 354, 547, 555 polishing with, 546–550 Q-switched, 342, 354, 386, 547 types of, 337 YAG, 348 Lawrence Livermore National Laboratory, 531 Lead as electrical contact, 600 LEED (see Film characterization) LIF (see Laser-induced fluorescence) Liquid immersion of substrates, 350– 353 Liquid-solid interface, 352 Loss tangent (see Dielectric) Low-energy electron diffraction (see Film characterization) Low-temperature, deposition, 408, 528, 529, 643 Lubricity, of diamond, 437
672 Mandrill (see also Substrate), 526 Manganese optical effects, 616 Mass spectroscopy, 88, 97, 131 Maxwell-Boltzmann distribution, 374 MCM (see Multichip module) Membranes (see Diaphragms) MESFET, 588 Metallization, of diamond, 553–557 Metal matrix (see Composites) Methane, 362 bending mode of, 394 emission from, 401 energy levels, 397, 404 hot bands, 400 stretch mode of, 393 Methyl radicals, 62, 64, 67, 70, 71, 87, 89, 93, 95, 100, 104, 106, 107, 109, 120, 130, 131, 134, 136, 146, 164, 214, 238, 242, 247, 268, 269, 271, 273, 287, 288, 296, 363, 407 Michigan State University, 227 Microcrystalline diamond, 456 Microcrystalline graphite, 343 Microporosity, in diamond, 454 Microstructure, of diamond, 444 interfacial, 463 Microsubstrates, Corp., 555 Microwave (see also Packaging) plasma enhanced assisted CVD, 5, 24, 56, 75, 92, 95, 99, 108, 126, 211–301, 528 power density, 279, 280, 292, 293 Microwave applicators, 215, 221 Microwave reactor types ASTeX, 229–231, 271 Berkeley, 231, 232 cavity, 227–229, 266, 277, 278, 280 France bell jar, 232, 233, 273 ellipsoidal, 233, 234 jet torch, 234, 235 magneto/ECR, 235–237 surface-wave, 234 tubular, 226, 227, 266, 269, 271, 273, 286
Index Mirage Effect (see Photothermal deflection) Mobility of holes in diamond, 581, 582 Modeling numerical, 285–296 thermal, of packaging, 558–561 Molecular beam mass spectroscopy, 88, 92, 100, 131, 271 Morphology, of diamond, 80–82, 86, 207, 223, 365, 452, 453 MOSFET, 591 MPD (see Multiphoton dissociation) Multichip module, 511, 512, 514, 521, 557, 562–569 Multiphoton absorption, 335 Mutual coherence function, 329 Nanocrystalline diamond, 268, 389 Nanoindentor, 633 National Institute for Research in Inorganic Materials (see NIRIM) National Institute of Science and Technology (NIST), 539 Navy, U.S., 629 Nickel as catalyst, 350, 352 as crucible, 351 as single crystalline substrate, 343 NIRIM, 24, 226 Nitrogen aggregates of, 613 effect on synthesis, 265 effect on thermal conductivity, 519 ESR active, 613 in polycrystalline diamond, 622 optical effects, 613, 621 NorCool, 569 Norton Diamond Film, 227, 448, 506, 569, 629 n-type diamond, 601 Nuclear particle dectors, 10 Nucleation, 37, 55, 73–79, 82–84, 110, 136 by abrasion, 83, 524 by carbon fibers, 525
Index [Nucleation] by fullerenes, 525 by metallic powder, 524 by substrate bias, 83–85, 136, 237, 238, 524 carbide layer, 78 density, 73, 74, 78, 82, 84, 86, 461, 523, 524 homogeneous, 74 incubation period, 75, 76 rates, 76, 77, 82 three-dimensional, 75, 76 ultrasonic, 524 OES (see Optical emission spectroscopy) Ohmic (see Contacts) Optical (see Absorption, Reflection, Substrates, Transmission) Optical absorption (see Transmission) Optical coatings, 10, 21, 355, 637–648 Optical emission spectroscopy, 92, 95, 96, 97, 136, 276, 278, 288 Optical materials, properties, 360 Optical pumping, 336, 337, 376–415 for diamond growth, 385–409 of liquids, 409–415 Optical reflection (see Reflection) Optical transmission (see Transmission) Optics, diamond, 607–654 Optoelectronic (see also Packaging) uses of diamond, 650–653 Overtone spectrum, 414 Oxidation of diamond, 438, 530, 636 Oxyacetylene, 340, 342 Oxygen growth, 260–265 optical effects, 624 Oxygen, in polycrystalline diamond, 622 Ozone, 336 Packaging (see also MCM) microwave, 521, 569 optoelectronic, 561, 562
673 [Packaging] radiofrequency, 521 three-dimensional, 521, 562–566 PCD tools, 427, 428 Permeable base transistor, 587 Phase diagram (see Carbon) Phase stability, 438 Phonon absorption multiphoton, 612, 623 single photon, 623 Phonon-phonon interactions, 344 Photochemistry, 359 Photoconductivity in diamond, 651–653 switching, 652, 653 Photodissociation, 362 Photolysis, 327, 339, 358–376 Photon detection in diamond, 650–652 emission from diamond, 653 infrared dissociation, 339 multiphoton dissociation, 336, 358 single photon absorption, 376 Photon-electron interactions, 344 Photon-phonon interactions, 344 Photothermal deflection, 538 radiometry, 537, 538 Physical Chemistry Institute, 23 Planarization, of diamond, 550–552 Plasma plume, 335 p-n junction, 580 Poisson’s ratio, 520, 530, 634 Polarization, 329 Polishing, of diamond chemical assisted, 545–546 for optics, 629, 630 ion beam, 549 laser, 546–550 mechanical, 542,543 thermochemical, 543–545 Polycarbonate, 496 Post deposition finishing, 447, 449 Precursors, for diamond, 364, 373–375 p-type diamond, 590, 594
674 Pyrolysis, 327, 339, 353 Pyrometry, 650 QQC process, 347–349, 353, 417 Quantum efficiency, 651 Quantum number, vibrational, 382, 396 Quantum yield, 362 Quench rate, 351 Radiometry (see Photothermal) Raman diamond peak, 340 D peak, 340, 341, 357 G peak, 340, 357 micro, 351, 357 quality factor, 626 spectra, 28, 180, 181, 190, 193, 203, 248, 340, 357, 358, 367, 394, 455, 456, 646 spectroscopy, 28, 148, 150, 343, 454 Rate constants for energy transfer, 411 of elementary reactions, 371 Raytheon, 555, 628 RBS (see Rutherford backscattering) Reactor scale-up, 98, 99, 119, 137, 141, 170–173 Recrystallized diamond, 351 Reflection, optical, 609 Refractive index of diamond, 8, 609 of several optical materials, 360 Sellmeier equation, 609 Relaxation time, 344, 345 REMPI (see Resonance-enhanced multi photon ionization) Resistivity, electrical of diamond, 7, 514, 519, 520, 530 Resonance defect, 383 Resonance-enhanced multiphoton ionization, 89, 90 Rf plasma CVD, 68, 98, 108, 141, 160, 171, 176, 198 Rhenium, 122, 128 Ring-down spectroscopy, 133 Rutherford backscattering, 36, 78, 127
Index Saint Gobain Industrial Ceramics, 448, 506 Sandia National Laboratories, 567 Saturation velocity, 580 SAW device (see Surface acoustic wave device) Scanning electron microscopy (see SEM) Scanning force microscopy (see SFM) Scanning tunneling microscopy (see STM) Scattering optical effects, 616–620, 626, 629, 639, 640 Schmidt number, 144 Schottky contacts, 584, 586 photodiodes, diamond, 652 Seals, diamond coated, 498 test results, 500 Secondary ion mass spectroscopy (see SIMS) Seeding (see Nucleation) Sellmeier equation (see Refractive index) SEM, 28, 30, 40, 78, 180, 182, 184, 201, 261, 267, 351, 356, 367, 393, 428, 448, 451, 453, 458, 467, 469, 470, 474, 476, 478, 493, 552 Semiconductor diamond as, 580, 598 Sensors, 12 SFM, 47, 48, 49 Shear modulus, diamond, 520 Silane, 339 Silica (see Substrates) Silicon (see Substrates) optical effects, 624 Silicon carbide diamond interface, 642 SIMS, 35, 36, 37, 591 Sooting, 109, 129 sp2 bonds, 4, 27, 28, 29, 71, 73, 170, 239, 242, 260, 263, 454 optical effects, 624, 625 sp3 bonds, 4, 28, 29, 73, 148, 239, 452
Index Spalling, 460 Specific energy, 141 Specific heat, of diamond, 7, 520 Stagnation flow, 347 Stainless steel, as substrate, 348 Stark-Einstein Law, 362 STM, 49, 50 Strain temperature of glass, 641 Strength of diamond burst strength, 520 compressive strength, 435 flexural strength, 633 mechanical strength, 519, 633–636 rupture strength, transverse, 435 shear strength, 530 tensile strength, 434, 530 Stress due to CTE mismatch, 437, 646–649 in films, 645–650 Substrate bias, 196–200 Substrate holder, 217–219 Substrate temperature, 85, 213, 214, 217, 218, 223, 240, 249, 250, 262 Substrates boron nitride, 75 diamond, 518–522 germanium, 355 glass, 641–650 optical, 637–650 silica, 641, 643 silicon, 78, 641, 642 tungsten carbide, 432 ZnS, 639 ZnSe, 639 Sun Microsystems, 531 Surface acoustic wave, 10, 12 Surface roughness, 447, 550, 629, 632 optical effects of (see Scattering) Tantalum, 122, 123, 128 TEM, 28, 30, 32, 33, 78 bright field, 340 of diamond, 340, 342 of graphite, 340
675 Temperature distribution, 345–347 gas, 275–278, 291, 292–294, 348, 385 gradient, 345 substrate, 85, 213, 214, 217, 218, 219, 223, 240 vibrational, 380 Tersoff potential, 63, 64 Textured growth, 37, 80, 81, 252–260, 631, 632 Thermal conductivity measurement methods, 535 of aluminum nitride, 514 of beryllium oxide, 514 of copper, 514 of diamond, 7, 9, 13, 43, 45, 46, 49, 435, 437, 514, 520, 530, 580, 636 of silver, 514 Thermal diffusivity, 45 Thermal expansion (see CTE) Thermal management, 11, 137, 141 Thermionic, 152, 155 Thermistors, 10 Thick film cutting tools, 430 Thin film cutting tools, 432 Thorium oxide, 152 3 method, 536, 537 Titanium as electrical contact, 589, 590 as filament material, 122, 128 Tools, cutting applications, 10, 137, 141, 441– 444 coated, thin film, 432 conventional, 426, 427 design of, 445–452 diamond adhesion to (see Adhesion) inserts, 475–478, 481–484 PCD, 427, 428 performance of, 441–445, 452–458, 475–498 thick film, 430–432 tungsten carbide, 426, 427 Tools, dressing, 501–503 Tools, rotary, 480, 481, 485–498
676 Transistors, diamond, 13 bipolar, 583 frequency, 582 history of, 583 IGFET, 589 MESFET, 582 MISFET, 589 MOSFET, 594–599 permeable base, 586, 587 point-contact, 584 power, 582 Translational energy, 327, 359 Transmission electron microscopy (see TEM) Transmission, optical diamond, 8, 9, 11, 520, 608 diamond coated substrates, 637–648 germanium windows, 355 of polycrystalline diamond, 616–621, 628 Treanor minimum, 381 Tungsten, as cathode material, 152 as electrical contact, 584 as filament, 58, 119, 122, 125, 527 as via fill, 556 melting point, 123 Tungsten carbide (see Tools) Turning (see Tools) Type I, 517 Type Ia diamond color of, 613 nitrogen in, 613 transmission spectra, 615 Type Ib diamond color of, 613 nitrogen in, 613 transmission spectra of, 614 Type II, 5, 7, 43 Type IIa diamond Cullinan diamond, 615 nitrogen in, 613 Type IIb diamond boron in, 613 Hope diamond, 615 transmission spectra of, 14
Index Ultraviolet (UV) (see also Absorption) diamond detector, 10, 12, 651 vacuum, 359 Union Carbide Corp., 19 Vacuum microelectronics, 12 Valence band, 610 Velocity, distribution, 374 Via, filling of, 555 Vibrational pooling, 389 Vibration electronic energy transfer, 377, 388 Vibration-rotation energy transfer, 404 Vibration-rotation-translation energy transfer, 340 Vibration-translation energy transfer, 340, 377, 382–384, 404 Vibration-vibration energy transfer, 377, 382–384, 404 Vicker’s indentation, 6, 648 Wavelength, 359, 608 Wavenumber, 359, 608 Wear resistance, of diamond, 439–441, 479, 504 Wear surfaces, 10 Windows, diamond endoscope, 627 gyrotron, 10, 11, 627 optical, 10, 11 Wire dies, 503–504 X-band, 582 XPS (see X-ray photoelectron spectroscopy) X-ray initiated reactions, 359 X-ray diffraction (XRD), 28, 36, 37, 40, 41, 78, 264, 457 X-ray masks, 11 X-ray photo electron spectroscopy, 88 X-RD (see X-ray diffraction) Young’s Modulus, 6, 9, 433, 434, 520, 530, 634 biaxial, 646 ZnS (see Substrates) ZnSe (see Substrates)
