Semiconductor diamond

C H A P T E R

2 Semiconductor diamond

C H A P T E R

2.1 Introduction Diamond has been regarded as the extreme wide bandgap semiconductor material in terms of its ultrawide bandgap, high thermal conductivity, and high electrical breakdown field as well as its high radiation resistance and mechanical and chemical robustness. Diamond is made up of single carbon atoms and has tetrahedral and covalent bonds between each carbon atom and its four closest neighbors. The lattice structure of diamond is formed by two vertical square lattices with one translation along the body diagonal line of the other. In the crystallographic cell of diamond, there are four carbon atoms in a face-centered cubic structure, and these four carbon atoms are located at one-fourth of the four diagonal spaces, respectively, as shown ˚ and the angle in Fig. 2.1.1. At room temperature, the C–C bond length of diamond is 1.55 A between the two C–C bonds is 109.28 degrees. Diamond belongs to the Fd3m space group. Naturally, diamond can be formed under high-pressure and high-temperature (HPHT) conditions, which exist more than 150 km beneath the Earth’s surface. In 1954, the US GE Corp. (GE) proved that they could successfully manufacture a synthetic diamond with the HPHT method (the ASEA Laboratory in Sweden also synthesized diamond in 1953, but it was not published). In 1968, Angus et al. applied low-temperature and low-pressure chemical vapor deposition (CVD) to successfully synthesize diamond films on natural diamond.

Ultra-wide Bandgap Semiconductor Materials https://doi.org/10.1016/B978-0-12-815468-7.00002-0

111

Copyright # 2019 Elsevier Inc. All rights reserved.

112

2. Semiconductor diamond

FIG. 2.1.1

The schematic of a diamond crystallographic cell.

However, the film quality was quite poor. In 1982, Kamo and Matsumoto synthesized microcrystalline diamond films on heterogeneous substrates by inventing microwave plasma CVD (MPCVD) and hot filament CVD (HFCVD), respectively. Since then, extensive worldwide research has been conducted by using CVD methods. In 1993, Jiang et al. prepared a (100) oriented diamond film consistent with the substrate orientation on the (100) oriented monocrystalline silicon substrate, which was called "highly oriented diamond film" with a very smooth surface. Since the early 2000s, the motivation to pursue electronic-grade diamonds has aroused the growth of single-crystal diamond (SCD) epilayers and wafers. Homoepitaxial growth techniques by using HPHT or CVD SCD substrates are often adopted for this purpose. Nevertheless, the limitation in SCD substrate size hampers the development of large (>1 in.) high-quality SCD wafers. To obtain larger SCD wafers, various techniques have been developed. These techniques include lateral expansion growth from the SCD substrate, clonal growth with a mosaic structure to enlarge the SCD wafer size, high-speed growth (>100 μm/h) of SCD epilayers by adding a small amount of nitrogen during CVD growth, and heteroepitaxial growth on an iridium substrate. Up to now, SCD wafers larger than 1 cm have been commercially available. A heteroepitaxial SCD wafer with a size of 92 mm was also reported. Intrinsic diamond is a perfect insulator with an ultrawide bandgap of 5.45 eV. Electrically conductive diamond can be achieved by doping impurities, usually during the gas phase deposition. Boron (B) and phosphorus (P) are the only acknowledged bulk acceptor and donor in diamond, respectively, up to now. However, both have high thermal activation energies (B: 0.37 eV, P: 0.57 eV). Interestingly, diamond owns a unique feature of surface conductivity with two-dimensional hole gas (2DHG) on a hydrogen-terminated surface, which provides a unique platform for field-effect transistors as well as low turn-on voltage vacuum electron emitters. As shown in Table 2.1.1, diamond has a high electron mobility of 4500cm2/V
s and a hole mobility of 3800cm2/V
s. It also has a high-saturation drift velocity (107 cm/s). The breakdown electric field (>10 MV/cm) is the highest among the wide-bandgap semiconductors. It has exceptionally high thermal conductivity (2400 W m1 K1), which is the highest among

113

HTHP synthetic diamonds

TABLE 2.1.1

Physical properties of diamond in comparison with other semiconductors

Materials

diamond

GaN

SiC

Si

Bandgap (eV)

5.45

3.27

3.4

1.12

>10

3.0

2.5

0.3

Drift velocity (10 cm/s)

1.5(e) 1.05(h)

2.0

1-2.5

1

Carrier mobility (cm2/V
s)

4500(e) 3800(h)

1000(e)

2000(e)

1450

Dielectric constant (εr)

5.37

9.7

8.9

11.7

Thermal conductivity (W/cm
K)

22

4.9

2

1.5

Johnson index (10 Ω
W/s )

2350

910

1080

2.3

Keyes index (10 W/K
s)

145

35

10

6.7

Baliga index (Si¼ 1)

43938

620

24

1

Breakdown field (MV/cm) 7

23

7

2

semiconductors. In addition, diamond is also the hardest material in nature and has excellent chemical inertness against acids. Due to these excellent properties, diamond has been extensively investigated as a new-generation semiconductor for applications in high-power, high-frequency, and hightemperature electronic devices, deep-ultraviolet (DUV) light emitting diodes, DUV photodetectors, and radiation detectors, especially for those applications in harsh environments. In this chapter, the state of the art for semiconductor diamonds will be reviewed. The topics include the crystal growth of diamond, impurities in diamond, electronic devices, DUV diodes, and DUV photodetectors.

C H A P T E R

2.2 HTHP synthetic diamonds Xiwei Wanga, Yan Penga, Shenglin Wangb, Dufu Wangb, Xiangang Xua a

State Key Laboratory of Crystal Growth, Shandong University, Jinan, P.R. China bJinan Diamond Technology Co., Ltd, Jinan, P.R. China

High-temperature, high-pressure (HTHP) growth is the first proposed method for labgrown diamond synthesis, taking full advantage of duplicating the natural diamond by

114

2. Semiconductor diamond

simulating the high-pressure and high-temperature conditions similar to those of the Earth’s underground. Meanwhile, the HTHP growth method, also known as the metal HTHP method or the temperature gradient HTHP method, creatively involves the combination of a metal alloy as the solvent catalyst with super high pressure and reasonable temperature aiming to accelerate the growth rate for the high-quality diamond single crystal [1–3]. By introducing diamond single-crystal seeds in the reaction cavity, the HTHP technique converts the graphite source to diamond under specific designed chemical-physical conditions. Nowadays, HTHP diamonds contribute much to the manufacturing industry, with wide use as a superhard application in abrasive powders, drill bits, and cutting tools. After decades of development, due to the decrease of equipment maintenance cost and the increase of the size and quality, HTHP diamonds also show themselves in the jewelry market and special optical applications. For the last 20 years, the technology of CVD diamond growth has attracted more and more attention. The best option for the substrate during the CVD single-crystal diamond growth is the high-quality HTHP diamond (100) surface with low impurity dislocation density [4, 5].

2.2.1 The creation of HTHP diamond synthesis 2.2.1.1 The former attempts In the late 18th century, scientists discovered that diamond was a pure material composed of only carbon [6, 7]. Many attempts were carried out to try to convert different kinds of carbon materials into diamond; however, all failed [8, 9]. Later, modern records show that Scottish chemist James Ballantyne Hannay was the first person who announced a synthetic diamond in 1880; however, it was proved to be a natural diamond later by a modern testing method [10, 11]. The second publisher was from France, chemist Henri Moissan, who won the 1906 Nobel Prize in Chemistry for his work on fluorine. Moissan used the advanced electric arc furnace at that time to prepare a synthesized diamond. After extensive research, he declared the discovery of the diamond-like mineral as silicon carbide. Finally, silicon carbide was named moissanite in honor of his work [12–15].

2.2.1.2 GE project for growing diamond General Electric, Carborundum Corp., and Norton joined together to develop the technology for diamond synthesis in 1941. Their piston-cylinder hydraulic pressure machine for the early study was able to heat carbon to around 3000°C under 3.5 GPa for several seconds. However, the work was interrupted due to World War II and resumed by GE’s Schenectady Laboratories in 1951. In order to raise the upper limit of the high pressure, Percy Williams Bridgeman, who won the Nobel Prize in Physics in 1946 for the study of physics under super high pressure, designed a new structure of the anvil using tungsten carbide, as shown in Fig. 2.2.1A. Later, Francis P. Bundy and H.M Strong made some improvements to the system design and successfully synthesized diamond. However, the pressure and temperature were difficult to control and the experiment failed to repeat [3, 17, 18]. In order to increase the maximum compression of the anvil edge, Tracy Hall designed the doughnut-shaped binding ring, which was called Belt in 1954. This successfully achieved the

115

2.2.1 The creation of HTHP diamond synthesis

COPPER COND‘T’R RING

BINDING BINDING RING RING

SOFT STEEL SAFETY RING 10 WONDERSTONE 4

3 NICKEL WONDERSTONE

SEMIPISTON

BINDING RING

BINDING RING

1

13

14

WONDERSTONE 8 STEEL 9

WC

RUBBER GASKET

SOFT STEEL SAFETY RING 15 16

WONDERSTONE 7 STEEL 5 6 NICKEL

2

Steel

COPPER COND‘T’R RING

11

12

COOLING WATER

∅ 12.7 ∅ 25.4

Catlinite gasket

AgCI pressure medium 0.25 mm

(A) FIG. 2.2.1

∅ 12.7 mm

(B)

(A) Bridgman anvil design for the 10 GPa pressure apparatuses. (B) Hall Belt press section structure [3, 16].

first lab-grown single-crystal diamond and this result was published in 1955. The invention of the Belt came up with the classic structure of the belt liked system design and HTHP diamond growth method. Fig. 2.2.1B shows the structure of the Belt system. A pyrophylite container was used in the Belt apparatus where the graphite and solvent catalyst were placed. The largest diamond diameter reached 0.15 mm, usable for industrial abrasive powder. After the successful replication of the work, Hall’s results were published in Nature, making him the first person in history to successfully synthesize a single-crystal diamond with a reproducible and well-documented process [16, 19]. Hall left GE in 1955 and joined Brigham Young University to become a professor of chemistry. He invented the multianvils machines “Tetrahedral” and “Cubic press” in order to avoid violating a US Department of Commerce secrecy order on the GE patent applications [20, 21] (Fig. 2.2.2). After publishing the results of the first HTHP lab-grown diamond, GE won admiration from the entire world. People believed a new age for diamonds was at hand, especially for the jewelry market. For more than 30 years, GE continued working on this project in the Belt machine and focused on the equilibrium state for the stable transformation from graphite to diamond. However, GE abandoned the jewelry market project due to the high cost as well as the small size and low repeatability for high-quality crystals. Several mechanisms were brought up by the researchers to describe the transformation from graphite to diamond. After summarizing years of experiments and productions, some of the researchers tended to believe that the transformation tended to be “Systems involving carbon dissolved in molten metals.” This theory was used to briefly explain the growth of the HTHP diamond for many years, which led to the rise of the famous “solvent-catalyst” mechanism principle [22, 23].

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2. Semiconductor diamond

(A)

(B)

FIG. 2.2.2 (A) Hall” Tetrahedral” designed high-pressure apparatus, (B) Hall “Cubic Press” designed highpressure apparatus [20, 21].

2.2.1.3 ASEA project for growing diamond Back in the 1950s, a Swedish electrical company named ASEA (Allm€anna Svenska Elektriska Aktiebolaget) in Stockholm also worked on diamond synthesis. In 1949, this company hired a team of five scientists and engineers and started a project named “Quintus.” In the 1980s, they claimed that they successfully synthesized a diamond one year before GE, and kept it secret for their business. It is worth mentioning that Baltzar Von Platen, who was the inventor of the first gas absorption refrigerator in 1922, was one of the team members. The basic principle for diamond synthesis was the same as GE. Von Platen’s team also knew that the processing needs high pressure and high temperature. For a long time, they were also trapped by problems in making a reliable machine strong enough for the required conditions. They finally designed diamond press equipment with six pyramid-shaped anvils pressing together to form a sphere around the sample. The whole structure was encased in a strong copper jacket and suspended in an alcohol-filled tank. Even like this, the whole process was quite dangerous and could have exploded when leakage appeared and led to security issues. The system usually broke during the operation and took days to repair. During their experiment, they also found that iron was the key to lowering the graphite melting point. Baltzar Von Platen left ASEA before the successful synthesis. According to the record from ASEA, on Feb. 16, 1953, nearly a year before the announcement from GE, their system ran under a pressure of 84,000 atmospheres at about 2000°C for an hour and the workers discovered small diamond crystal powders after unwrapping the carbon parcel. However, they decided to cover the research result in case any competitors stole from other companies. As a result, the world never officially recognized their work as the first successful synthetic diamond [24, 25].

2.2.2 The HTHP diamond press apparatuses

117

2.2.1.4 Shock wave method for HTHP diamond In 1962, the Stanford Research Institute reported the explosive shock apparatus for smallsized diamond synthesis. The system is detonated against a free piston to direct a planar shock wave into a solid cylindrical specimen. According to the report, the wave generated in approximately a microsecond and was determined to be of the order of 400–500 kBar and about 1000–1500°C. The considerable graphite had been successfully converted directly to diamond in this process. This method is now still used for the nanometer-sized fine abrasive powders [26, 27].

2.2.2 The HTHP diamond press apparatuses Now, overall in the world, the HTHP diamond press apparatuses are all solid-state static equipment, the maximum limit temperature reaches about 2500°C, and the pressure is up to 100 GPa. The system ran into stable pressure and temperature conditions and lasted from hours to days for diamond crystal growth. The basic principle for building up high pressure is based on one simple theory: that decreasing volume leads to increasing pressure. However, this process has cracking problems, even for the hardest materials. Stricter requirements call for higher temperature and system duration in the growth stage. As a result, the evaluation of the HTHP diamond growth apparatus finally focus on the higher processing configuration of temperature, pressure and service life. The first successful creation was a well-known opposed anvil system named Belt, as mentioned above. However, the most serious problem for Belt is the lifetime of the anvil materials (Fig. 2.2.3). The damage of even one anvil marge in the high-pressure stage brought a dangerous explosion and led to a rapid pressure drop, which led to a failure of the diamond crystal growth and also high probability to destroy the opposite anvil and other parts of the equipment. This increases the cost for diamond synthesis and has security issues. With decades of development for Belt, the system is much more stable compared to the prototype and is still used in Central Europe, Japan, and Russia as well as colleges and institutes [20, 28, 29].

Press platten

Ceramic tube

Heater sleeve

Tungsten carbide anvil

Steel support rings for the anvil

Tungsten carbide die

Steel support rings for the die

Gaskets

Reaction volume

Press platten

Electrical contact

(A) FIG. 2.2.3

(B) (A) Schematic diagram of the Belt apparatus. (B) Steel support rings for the die for Belt [19].

118

2. Semiconductor diamond

2

3

1

sizes not in scale

Anvils

4

Synthesis capsule Pyrophyllite

innner (x6) graphite tungsten outer (x8) steel sizes not in scale

Inner anvils Outer anvils Rubber diaphragm Hydraulic oil Pressure vessel

R 500 mm

(A)

(B) FIG. 2.2.4

(A) Schematic structure of BARS. (B) Photo of BARS [30].

Fig. 2.2.4. is the “BARS” technology invented during 1989–1991 by scientists from the Institute of Geology and Geophysics of the Siberian Branch of the Academy of Sciences of the USSR, which is also known as the “split sphere.” Nowadays, the typical press apparatuses are able to reach 10 GPa and 2500°C. A 2-cm3 specially designed ceramic cylindrical reaction cavity can be placed in the center of the device with cubic-shaped material for pressure transmitting and temperature gradient distribution. Eight anvils around the cavity are assembled to form a sphere shape in order to provide pressure from different directions. Due to the effective pressure stage during growth, BARS is well known for the synthesis of large single-crystal diamonds, up to more than 10 ct for one piece; the growth rate is able to reach 20 mg/h [30–32]. The “cubic press” is now the most popular equipment for the mass production of industrial diamonds. It is a type of multianvil deformation apparatus that uses six cubically arranged anvils to provide independent pressurization and deformation of the sample. Four anvils are oriented

119

2.2.3 Theory of the HTHP diamond growth

Cubic sample cell

(A) FIG. 2.2.5

(B)

Anvil

(A) The photo of the entire cubic press system. (B) Schematic diagram of the cubic press [33].

in the horizontal opposing at 90 degrees and the remaining two are oriented in the vertical to achieve a symmetrical cubic interspace for the reaction cavity, which is also known as the “synthesis cubic.” The system is equipped with a hydraulic pump station providing stable pressure during the HTHP stage as well as an electrical heating system and water cooling system for the anvils to achieve a steady-state temperature gradient during diamond growth. The diameter of the synthesis cubic reaches 50–80 mm; it is available to introduce more seeds into the reaction rather than one in the BARS and Belt processes. Taking full advantage of production quantity with lower cost, cubic press is now producing more than 90% of the single-crystal diamonds in abrasive, jewelry, and superhard applications every year [33, 34] (Fig. 2.2.5).

2.2.3 Theory of the HTHP diamond growth 2.2.3.1 The thermodynamics of the graphite to diamond transformation Because both diamond and graphite are different phases formed by an elementary carbon substance, the formation of diamond from graphite is simply a phase transformation, which is described as follows: Cgraphite ! Cdiamond

(2.2.1)

Assuming the reactions happen in the sealing system with an isothermal pressure and temperature, the Gibbs energy changes for the transformation reaction can be described as:

120

2. Semiconductor diamond 0 Δr G0m ¼ Δr Hm  TΔr S0m

(2.2.2)

Where ΔrG0m, ΔrH0m, and ΔrS0m are the changes for the Gibbs free energy, enthalpy, and entropy, separately. At room temperature and atmospheric pressure, the enthalpies for diamond and graphite are 0 and 1.8962 KJ/mol, respectively, and the entropies are 5.6940 J/(K
mol) and 2.4389 J/ (K
mol), respectively. As a result, we can calculate Eq. (2.2.2) as, 0 ¼ 1:8962ðkJ=molÞ ΔH298

ΔS0298 ¼ 3:2551ðJ=ðK
molÞÞ

(2.2.3)

0 Δr G0298 ¼ Δr H298  TΔr S0298 ¼ 2:866ðkJ=molÞ

Normally, besides the room temperature, graphite is the stable phase at any temperature in the atmospheric pressure, which means that the Gibb’s energy changes from graphite to diamond are always quantitatively larger than zero. The thermodynamics analysis result proves that the transformation would never react without the extra work from the surroundings. However, because diamond is in a more dense form, more energy is expected for its formation [34–37]. The direct form of the energy transformation is temperature and pressure. When the pressure and temperature are high enough, the reaction should finally come into an equilibrium, keeping the transformation in a stable process. Therefore, the expression for the Gibb’s energy changes of the chemical potential Δ Gp1, T0 is given by the equation: pð1

ΔG0p1 , T  ΔG0p2 ,T

¼

ΔVR dP

(2.2.4)

p2

Where △ Gp2, T0 is the Gibbs energy change at a pressure of p2 and temperature T.At high pressures and temperatures, the volume change is a function of both T and P, which cannot be assumed as a constant. The expression is given in Eq. (2.2.5) to calculate the Gibbs energy changes at any temperature and pressure. The density of the diamond and graphite are 3.513  103 and 2.261 103 kg/m3, respectively. Then, Eq. (2.2.4) can be simplified as follows: ΔGp, T ¼ 1:8962 + 0:0032551T  1:89559  106 ðp  101:325Þ kJ=mol

(2.2.5)

When the pressure is under 1GPa, the temperature ranges from 298 to 2273 K. In this case, the Gibbs energy changes are all smaller than zero, which means graphite is the stable phase at any temperature in the atmospheric pressure. By increasing the pressure (above 2 GPa), the Gibbs energy changes to the negative, which leads the transformation from graphite to diamond to react spontaneously. Meanwhile, the diamond phase becomes more stable (Table 2.2.1) [38].

2.2.3.2 The kinetics of graphite to diamond transformations It is a general belief from the thermodynamics theory that diamond is the metastable state and graphite is the stable phase. However, at atmospheric pressure and room temperature, diamond seems not to tend toward transformation into graphite itself due to the high energy barrier that is necessary for breaking the sp2 covalent bond between the graphite carbon atoms. The energy

121

2.2.3 Theory of the HTHP diamond growth

TABLE 2.2.1

Shows the calculation for the Gibbs energy changes as temperature and pressure Temperature(K)

ΔG (kJ/mol) Pressure (kPa)

298

773

1023

1273

1773

1973

2273

101.325

2.8662

4.4123

5.2261

6.0398

7.6673

8.3183

9.2948

10000

2.8474

4.3936

5.2073

6.0211

7.6486

8.2996

9.2761

100000

2.6768

4.2229

5.0367

5.8504

7.4779

8.1289

9.1054

1000000

0.9708

2.5169

3.3307

4.1444

5.7719

6.4229

7.3994

2000000

0.9248

0.6213

1.4351

2.2488

3.8763

4.5273

5.5038

4000000

4.7160

3.1699

2.3561

1.5424

0.0851

0.7361

1.7126

6000000

8.5072

6.9610

6.1473

5.3335

3.7060

3.0550

2.0785

required for the transformation is the reaction activation energy that relates to the probability for the carbon atoms to step over the barrier after absorbing enough energy [39]. As a result, the higher reaction activation energy leads to the higher reaction rate. For most of the chemical reaction, the most effective way to increase the reaction rate is increasing the temperature. The Arrhenius equation is usually used to describe the temperature dependence of reaction rates. Ea

k ¼ Ae RT

(2.2.6)

where k is the rate constant, T is the absolute temperature, A is the preexponential factor for each individual chemical reaction, and R is the universal gas constant. The Arrhenius equation demonstrates the relationship that reaction rate k is in positive correlation with the temperature T, and the calculation proves for most of the chemical reaction that k increases about 2–10 times when the temperature increases 10°C [40].

2.2.3.3 The driving force for diamond growth The metal transfer for the HTHP method is driven by the oversaturation of carbon, which is affected by the temperature and concentration gradient in the cavity. The relationship for the chemical potential (μ) with the entropy (s) and pressure (v) are as follows, dμ ¼ sdT + vdP

(2.2.7)

dΔμ ¼ ΔsdT + ΔvdP

(2.2.8)

When the temperature difference △ T ¼ 0, Eq. (X) can be transformed to,   Δμ ¼ μg  μd ¼ RT ln Xg =Xd ¼ ΔvðP  Pe Þ

(2.2.9)

ΔX=Xd ¼ ðΔv=RT ÞðP  Pe Þ

(2.2.10)

ΔX ¼ Xd  Xg

(2.2.11)

Xd and Xg are the solubility of the diamond and graphite in the reaction [41, 42].

122

2. Semiconductor diamond

FIG. 2.2.6 Schematic diagram of the structure in the reaction cavity [33].

Temperature gradient 20–50°C Graphite

Metal melt High temperature

Diamond crystal Low Growing crystal seed temperature

During diamond growth, outer shell intermediate metal alloy carbide films of about 100 micrometer thickness exist around the crystal, providing a channel for the graphite to transform into diamond. The temperature changes little through the film and the driven force for the diamond growth is the supersaturation for the carbon sources [43–45]. The relationship between the chemical potential changes with temperature is as follows,  d d =XL ¼ ΔSf ΔT (2.2.12) Δμ ¼ μH  μL ¼ RT ln XH   ΔX=Xd ¼ ΔSf =RT ΔT (2.2.13)   ΔX=Xd ¼ Δhf =RTTm ΔT (2.2.14) XdH and XdL correspond to the solubility for the diamond in high and low temperature areas, respectively and μH and μL are the chemical potentials. △ Sf and △ hfare the dissolving entropy and enthalpy, respectively, and △ T ¼ TH  TL,△X ¼ XdH  XdL. Tm is the melting point of diamond. The whole cavity is in the accurately designed temperature gradient field, which leads to carbon transfer from the high-temperature area into the low-temperature area. The diameter of the cavity is less than 10 mm. By accurately controlling the electricity heating power and cooling water flow both in the upward and downward anvil, the temperature gradient take place. The graphite melts into the metal alloy in the high-temperature region and the seeds are in the low-temperature region. The temperature drives the carbon flax to transmit from the metal alloy downward to the seeds and deposit systematically during diamond growth [33, 46]. Fig. 2.2.6 shows the structure for the inside core of the cubic synthesis; the temperature difference of the growth space for the diamond crystal is only 20–50°C and the ideal gradient difference should be limited to 20–30°C.

2.2.5 Metal alloy solvent-catalytic theory

FIG. 2.2.7 Phase diagram of the carbon element, which provides a short description to understand the phase transfer constructed from the experimental data and calculations [47].

120

30

Pressure [GPa]

123

Shock wave synthesis

Liquid carbon

25

Diamond 20 15

10

HP-HT synthesis Catalytic HP-HT synthesis

Graphite

5 0

0 1000 CVD diamond deposition

2000

3000

4000

5000

Temperature [K]

2.2.4 Phase Diagram of the HTHP diamond synthesis Fig. 2.2.7 shows the phase diagram of the carbon element. The solid lines in the diagram show where the coexistence between the two adjacent phases is. The line between the diamond and graphite phase shows at what values of the P and T qualitatively for both of the two phases are in equilibrium with each other. Diamond growth conditions can be classified as four regions for the phase diagram. The white region is the stable phase for the HTHP diamond while the black is graphite as the stable phase instead. The green region is the shockwave method under a sudden enhanced super high pressure for the small grains of synthetic diamond; the hexagonal diamond phase may exist in this region. In the gray region, liquid carbon is more stable where the diamond can be easily graphitized. Without the assistance of the metal alloy catalyst, the transformation from graphite to diamond has little chance to react according to the phase diagram because the temperature and pressure requirements almost double (from the temperature and pressure of the red region to the purple region). The preparation cost is high and the process is difficult to achieve a large, highquality synthetic diamond single-crystal growth [47–50].

2.2.5 Metal alloy solvent-catalytic theory One of the most important creative discoveries of the HTHP diamond synthesis was to introduce a metal alloy into the reaction system for the diamond crystal growth. The increasing reactivity from graphite to diamond is affected by the solubility of the carbon in the metal. For decades, researchers kept exploring the compound formula of the metal alloy for certain temperature and pressure. Table 2.2.2 shows some metal alloy compound formulas used as the solvent catalyst for differently colored jewelry applications. Researchers tried different temperatures and pressures for the synthesis and listed the minimum requirement for diamond growth. The metal alloy dissolves graphite at elevated pressures ranging from 5 to 10 GPa when heated to 1300°C to 1500°C, this stable growth process is proved to be a technical and commercial reasonable for the apparatus generation evolution during the latest half century. With the

124

2. Semiconductor diamond

TABLE 2.2.2

Large HTHP diamond grown for research purposes [51] Growth rate (mg h21)

Nitrogen content (ppm)

Morphology

Near colorless 300

2.3

<0.01–2

{111}, {100}>{113}, {115}.(110}

4.6

Near colorless 500

1.8

0.4–1

{111}>{113}>{100}>{110}

Fe-Co

25

Yellow

1000

5

100–1000

{111},{100}»{110}, {113}

Fe-Al-B

5.1

Blue

760

1.3

–

{111}.{110}.{100}, {113}, {115}

Solvent/ catalyst

Maximum size (ct)

Color

Co-Ti

3.4

Fe-Al

Growth time (h)

help of the accurately controlled temperature gradient and diamond seed, the graphite resource powder recrystallizes into a large diamond crystal in the reaction cavity [51]. Several principles were brought up to explain the application for the metal alloy in the transformation reaction during the HPHT stage. Among all, the most general believed model is the “solvent-catalytic theory.” This theory considers the HTHP method as a solution growth that the metal alloy works as a solvent to the carbon during the transformation and also as a catalytic to reduce the requirement for the temperature and pressure [5, 23, 36, 41, 48]. When the diamond growth begins, the molten metal starts a eutectoid transformation combining the carbon atoms and forms a shell film of around 100 micrometers covering outside the diamond crystal surface after the nuclei process. This plays an important role as a transmission intermediate for the carbon atoms to transfer from graphite and separate out to the diamond interface. Several reactions take place simultaneously in the cavity during growth, according to the solvent/catalytic theory. The process is as follows:

2.2.5.1 The graphite carbon dissolves into the metal alloy The temperature and concentration gradient act as a driving force for the graphite to contact and spread into the molten metal alloy. Most of the reactions take place in the interface and pass through the metal, which is acting as a solvent. The solubility of the carbon is one of the most important characteristics for the transformation [23, 51].

2.2.5.2 Chemical reaction transfers graphite carbon to diamond carbon The secret for the metal alloy element selection is in the special outer electronic structure and crystal lattice constant. By employing the metal alloy into the reaction cavity, the energy requirement for the diamond synthesis overcomes the energy maximum barrier; the total transformation represents an activated stage for the reaction. When using the Fe-based catalyst as the solvent-catalyst, both EELS and X-ray diffraction analysis proved the existence of the micrometer thickness of γ-Fe and other Fe carbide forms such as Fe3C in the outer shell film of the metal alloy after the growth. The Bin Xu team compared the carbon and Fe EELS spectra for the different positions of the outer shell of the Fe-Ni-C catalyst and concluded that the existence of a pearlite Fe3C (maybe also Ni3C, which is not stable in the atmospheric pressure and room temperature; however, it may exist during the high-temperature, high-pressure stage) structure provided the transformation from the sp2 carbon bond to the sp3 bond [45, 52–55].

125

2.2.6 The study in the metal alloy element

2.2.5.3 Sp3-bonded carbon atoms decomposed from Fe3C and deposition on the diamond surface Finally, the sp3 band carbon resolves from the interface and builds the structure of the diamond phase. The existence of austenite γ-Fe accelerates the decomposition of the Fe3C to iron and carbon atoms while the iron atoms are able to generate Fe3C and γ-Fe again to repeat the entire reaction. In all processes of carbon transformation, atomic rearrangement takes place. The bonds between the iron and carbon atoms break and the carbon atoms separate out due to the oversaturation to form new sp3 bonds with the surrounding carbon atoms, arranging in regular periodic arrays and displacing to a new equilibrium position for the transformation [56–58]. However, employing in situ analysis in the closed cavity during the HPHT growth for the real-time process is technically difficult. The general method to study the reaction in the cavity is to understand the metal and carbon material property after the diamond crystal growth at the normal temperature and pressure. This makes it difficult to predict the true mechanism during the growth stage. More experimental studies are needed to understand the mechanism.

2.2.6 The study in the metal alloy element Table 2.2.3 shows commonly used elements for the metal alloys and a comparison of the outer electron configuration and the number of unpaired d shell electrons. The way to increase reactivity is to increase the number of d vacancies in the electric orbital due to the increasing reaction with the carbon atoms. The solvent-catalyst transition elements share TABLE 2.2.3

Data for elements used for industry metal alloy [59]

Without unpaired d electrons

With unpaired d electrons

Element

Electron configuration

No. of unpaired d-shell electrons

Element

Electron configuration

No. of unpaired d-shell electrons

Zn Zinc

3d104s2

0

Ce

4f15d16s2

1

Mg Magnesium 3 s Al Ge

Aluminum

2

0

1

0

3p

2

2

Germanium 4s 4p

0

Ag Silver

10

1

4d 5s

Au Gold

l0

1

10

1

Cu Copper Si

Silicon

5d 6s

3d 4s 2

1

3s 3p

0 0 0 0

La

Cerium

1

2

5d 6s

1

5

2

3d 4s

5

Nickel

8

2

3d 4s

2

Cobalt

7

2

3d 4s

3

6

2

3d 4s

4

2

2

3d 4s

2

Chromium

5

1

3d 4s

5

Vanadium

3

2

3d 4s

3

5

1

5

Tungsten

4

2

5d 6s

4

Platinum

9

1

1

Lanthanum

Mn Manganese Ni Co Fe Ti Cr V

Iron Titanium

Mo Molybdenum 4d 5s W Pt

5d 6s

126

2. Semiconductor diamond

the same prosperity that the atoms are all lack of 3d orbital electrons, which helps them to attract 2Pz electrons of the carbon atoms during the HPHT stage and transforms the carbon electron configuration from graphite sp2 hybrid bond into the diamond sp3 state. The nearby carbon atoms are attracted to the phase change gradually and result in an increase of the carbide compound melting point to generate a more reasonable carbon-metal ratio, which leads to the acceleration of the entire HTHP growth rate [59]. In order to increase the carbon solubility, metal elements with more d orbital electrons shall be chosen. However, what the periodic table shows is the electron configuration at the room temperature and atmospheric pressure. It is believed that more unpaired d orbital vacancies bring the high possibility of bonding with the carbon atoms for the transformation. While in the HPHT stage, electrons tend to move to a more active energy orbital. For most transition metal elements, the d orbital energy is higher than that of the s orbital in a higher quantum number. The valence electrons shall first fill in the latter. However, this rule changes when it comes to an element with five electrons, acting as either half or complete fully in the d orbitals. Take Cr and Mo for example, the electron will move to one or two into the d orbitals from the s orbital for a more stable form. Palladium (Pd) is another exception for the fully occupied 10d orbitals (10 electrons); the configuration makes it difficult to by active reaction to be a catalyst element. However, the energy for the fully occupied electrons is only slightly lower than the 5s orbital electrons. During the growth stage, the 4d electrons can be easily moved into the 5s orbital, leaving a vacancy for the interaction with a carbon atom. In contrast, elements such as Cu (3d104s1), Ag (4d105s1), and Au (5d106s1) have already fully occupied the d orbital, providing one extra electron into the orbital of the higher quantum number. These elements have poor ability in dissolving graphite due to the inertness of the d orbital electrons in the high-temperature, high-pressure stage [39, 59, 60]. Wakatsuki et al. reported that elements such as Cu, Ag, and Au may also show a solvent catalytic effect with a combination of other elements. They separate the alloy element into a carbon attractor and carbon repeller, which act as the carbon receiver and donor, respectively. The carbon attractor includes group IVa, Va, VIa element with unpaired d electron vacancy (Mo, W, Ti, Zr, Hf, V) while carbon repeller include group Ib (Cu, Ag, Au). The combination of both the two group elements may also present a solvent catalyst character for diamond growth, but this is not effective compared to the single element catalyst such as Fe. This theory was later questioned by Bundy, who suspected that the copper was contaminated by other catalytic metals. Kanda et al. achieved diamond growth with a Cu catalyst at 60 Kbar and 1600°C [23, 59–61]. Actually, the catalytic ability for diamond growth can be measured by the active energy provided by the temperature and pressure. Experiments proved that the noncatalyst material could be transformed into a catalyst material at high temperatures. Even nonmetals such as P, CaSO4, or CaCO3 are able to achieve diamond growth at a threshold temperature (Table 2.2.4) [[62, 63]]. Another theory for choosing the solvent catalyst elements is the vertically aligned and lattice pucker principle by Gou in China. Diamond density is only about 56% that of graphite. The total volume shall decrease when the transformation actives and the lattice of the graphite shall be puckered during the reaction. For example, the (0001) planes of the graphite shall be puckered into diamond (111) planes. Besides the difference of the electronic configuration, the size of the solvent catalyst atom is another important point that systematically changes the lattice angle and length. One single atom is not able to make a difference for the diamond

127

2.2.6 The study in the metal alloy element

TABLE 2.2.4

Minimum conditions for diamond growth [39]

Catalyst

P (Kb)

T (°C)

Ni(80)–Cr(14)–Fe(6)

45

1150

Mn(92)–Cu(8)

48

1400

Fe(67)–Ni(33)

50

1280

Co

48

1450

Mn(92)–Co(8)

50

1450

Ni–Cr

51

1450

Fe

51

1400

Mn–Ni

53

1475

Mn(92)–Ni(8)

53

1475

Ni

52

1400

Mn

54

1500

Pt–Co(25)

55

1500

Pt(80)–Co(20)

55

1500

Rh

57

1700

Cu

60

1600

Ta

60

1880

Pt

63

2000

Cr

63

2100

CaSO4

77

1700

CaCO3

77

1800

P

77

1800

phase conversion, instead, a mass of them matters. However, the metal alloy lattice atoms acting with the whole surface of the graphite lattice shall be much more efficient. Therefore, it is important to form a bonding interface with every metal atom vertically aligned with the carbon atoms bonding together. Gou pointed out that the edge length of the equilateral triangle form shall be close to that of the three diamond atoms, forming a vertically aligned array for the bonds between metal atoms and carbon atoms. For the metal lattice, the atom edge lengths for a face-centered cubic structure (100) surface, a diamond-type structure (111) sur˚ (graphite) face, and a closed-packed hexagonal structure (0001) surface shall be around 2.46 A ˚ and 2.51 A (diamond) in order to have a quantity of chemical bonds for the metal alloy with the carbon interface. These chemical bonds are able to drag the carbon atom away from the graphite interface and puck the structure from graphite into diamond. His model explains the orientation relationship in the interface and the existence of the Fe3C after diamond growth [58, 64–66].

128

2. Semiconductor diamond

2.2.7 HTHP single-crystal diamond type and color centers The diamond types are classified by the level and type of chemical impurities. The diamond is normally separated into four types due to the impurity of different elements and the color center in the lattice: Type Ia, IIa, Ib and IIb. The impurities measured are at the atomic level within the crystal lattice of carbon atoms. Different types can coexist in a single crystal, as a natural diamond often mixes with Type Ia and Ib. Table 2.2.5 shows the differences of diamond type [67]. Being a neighbor of carbon in the periodic table, nitrogen is the main impurity for both natural and HTHP diamonds in single and aggregated forms. Most HTHP diamonds for industry applications contain a high level of nitrogen in the A and C nitrogen center form, as type Ia and Ib, which correspond to the electrical neutral pairs and single atoms in the neighboring substitutional position in the diamond lattice, respectively. The N+ center is also regular in the HTHP diamond lattice, which responds to the single substitutional nitrogen in charge state +1. The N+ center is believed to exist due to the impurity of Ni ions in the metal alloy. The C center imparts a yellow or brown color due to the different concentrations of isolated nitrogen impurity while the A center is nearly colorless [68, 69]. HTHP diamonds absorb both infrared and ultraviolet light. Khokhryakov et al. used FTIR to map the (110) plates of the HTHP diamond. It was indicated that the nitrogen for the areas adjacent to the seed was mainly in the form of A-center and the peripheral areas were dominated with C-center, leading to an apparent optical transmission in color for the visible spectrum range. The features of the spectrum are the C-center with the main absorption at 1130 and 1344 cm1. Babich et al. studied the distribution of the color center of the HTHP diamond grown in an Fe-Ni-C metal alloy system also by FTIR mapping. They tried to explore the H1a center in the HTHP diamond with an IR peak at 1450 cm1, which is related to the di-nitrogen interstitial in the diamond lattice with a complex of two nitrogen atoms sharing a single site in a (001) split configuration. The FTIR mapping data showed the distributions of the C, A, N+, and H1a center. The concentration of the C center increased in the growth directions from the seed to the outside parts of the octahedral growth sectors while the N+ center decreased simultaneously. A center was in the inner and middle areas of octahedral growth sectors, absent in the vicinity of the outer growth surface, and the H1a center was in the adjunction area for the A and C centers. It was concluded that the H1a center was only associated with the transition zone where nitrogen was partly transformed from the C center into the A center during the growth of the HTHP diamond, playing a role as some intermediate state in a step-wise reaction of transformation from the C to A center [70, 71]. TABLE 2.2.5

Diamond type

Diamond type

Ia

Ib

Natural content ratio

98%

Nitrogen impurity, ppm

210

0.1% 3

110

3

IIa

IIb

1 2%

0%

1

1 100

Boron impurity, ppm Color

Colorless-Yellow

Yellow-Brown

Colorless

Blue-Gray

129

2.2.7 HTHP single-crystal diamond type and color centers

It is also possible to control the crystal optical color with different compound formulae of metal alloys. In order to grow a colorless IIa diamond type for jewelry, aluminum, zirconium, and titanium are the most common elements for the “nitrogen getter” during the HTHP stage. These getter atoms trap the nitrogen from the atmosphere or the graphite source to stop nitrogen dispersal into the diamond lattice by forming stable nitrides and carbides. The final growth rates for the high quality IIa diamond are significantly lower than type Ib. A large colorless diamond requires necessitating a longer growth stage and more effective control of the temperature and pressure [73]. Burns et al. found that the nitrogen getter efficiency decreased with increasing temperature because of the difference between the free energies of formation of the respective carbide and nitrides decreases, as shows in Fig. 2.2.8. The nitrogen concentration in the diamond increases as the temperature increased from 1600°C to 1900°C [72]. Boron is another element next to carbon in the periodic table that can be incorporated into the diamond lattice due to a similar atom size. By adding boron in the graphite source or metal alloy, the diamond crystal appears to be blue or gray in color and the type changes to IIb. The absorption spectrum of boron makes the diamond crystal in type IIb absorb the red to yellow color. These diamonds are a p-type wide-gap semiconductor material with high carrier mobility, high thermal conduction, and extreme elastic properties. The thermal activation energy of boron in diamond is around 0.37 eV. Even by carefully controlling the compounds of metal alloys, carbon flux, and temperature gradient, inclusions from the metal alloy such as Fe and Ni are able to incorporate into the growing diamond. These inclusions often bring large stress to the crystal and appear as a rod shape surrounded with feathers or fractures. Some of the crystal may even be attracted to a magnet. Numerous Ni-related defects have been detected. The major form of the structures is distinguished as substitutional Ni, nickel vacancy, and nickel vacancy complex decorated by substitutional nitrogen [76]. Fig. 2.2.9 shows the different sharpnesses of the morphology of the HTHP diamonds (The authors thank Mr. Dufu Wang from Jinan Diamond Technology Co., Ltd for providing sample

Nitrogen concentration (ppm)

8

6

4

2

0 1600

1700

1800

1900

Synthesis temperature (˚C) FIG. 2.2.8

The dependence of nitrogen concentration in an HTHP diamond with aluminum in a metal alloy as a function of growth temperature [72].

130

(A) FIG. 2.2.9

2. Semiconductor diamond

(B) (A) Photograph of natural diamonds and (B) HTHP diamonds.

photos. The morphology of the crystal depends on the ratio between the lateral growth rate and the vertical growth rate. (100) and (111) are the main growth habit planes for the HTHP diamond. The growth coefficient rateα and β parameters of the HTHP diamond are defined as the growth rate ratio of the growth rate of the (100)/(111) and (100)/(110) surface.   pffiffiffi V100 α¼ 3 V111   pffiffiffi V100 β¼ 2 V110 The variables of V100, V111, and V110 are the growth rates in the [100,111,110] direction, respectively. Normally, α is used to describe the sharp morphology of the crystal from the outer parameter. When α ¼ 1, the crystal appears to be in cubic morphology while the diamond shows an octahedral structure shape when α ¼ 2.85. For natural and HTHP diamonds, the growth rate ratio ranges from 1 to 2.85. The morphology of the natural high-quality diamond usually appears to be octahedral or dodecahedral in shape while HTHP diamonds in all type tend to be cubo-octahedral. The present facets are (100), (111), and (110), usually present as cubic faces, octahedron faces, and rhombic dodecahedron faces [77, 78]. Fig. 2.2.10 shows the phase diagram of the Ib and IIa diamond. Generally, at the same pressure, the cubic (100) growth sectors predominate at the lower temperature, whereas (111) is dominant at higher pressure. The temperature span for areas A, B, and C at 5.5 GPa is about 40°C, 20°C and 10°C. As a result, precision temperature control is the key point for morphology control [79, 80].

2.2.8 Summary The historical milestone of the HTHP diamond as well as some aspects of the issues related to the synthesis and characterization of the HTHP diamond were reviewed. The last century

131

References

A

B

{100}

C

{100}+{111} (skeleton)

6 {111} (Inclusion)

5.5GPa

5

Diamond/Graphite equilibrium line

Pressure (Gpa)

Pressure (Gpa)

(skeleton)

6 {111}

5.5GPa

5

(A)

1350 1400 Temperature (˚C)

(Inclusion)

Diamond/Graphite equilibrium line

Solvent/Carbon eutectic melting line

Solvent/Carbon eutectic melting line 1300

{100} {100}+{111}

1300 (B)

1350 1400 Temperature (˚C)

FIG. 2.2.10 (A) Growth region of type Ib diamond. (B) Growth region of type IIa diamond [75].

has witnessed the development of HTHP diamond growth, which is a fruitful multidisciplinary combination of crystallography, thermodynamics, high-pressure physics, and mechanical design. Driven by the motivation of industry machining and jewelry applications, the HTHP diamond has been a superstar since the first successful synthesis. However, as with all other crystal growth, the main obstacle remains the limited growth rate, which renders the process and the product expensive. Furthermore, the growth of the high-quality crystal is also an important issue. There are still many questions in theory and mechanism of the HTHP crystal growth, metal alloy solvent catalyst, and high-pressure transformation to be further answered.

Acknowledgments The authors would like to acknowledge Mr. Dufu Wang, Mr. Shenglin Wang and Mr. Changjiang Liu from Jinan Diamond Technology Co.,Ltd for providing significant contributions to the writing of this chapter as well as the HTHP diamond photos.

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[71] Y.V. Babich, B.N. Feigelson, D. Fisher, A.P. Yelisseyer, V.A. Nadolinny, J.M. Backer, The growth rate effect on the nitrogen aggregation in HTHP grown synthetic diamond, Diam. Relat. Mater. 9 (2000) 893–896. [72] R.C. Burns, J.O. Hansen, R.A. Spits, M. Sibanda, C.M. Welbourn, D.L. Welch, Growth of high purity large synthetic diamond crystals, Diam. Relat. Mater. 1433 (1999) 8. [73] I. Kiflawi, H. Kanda, S.C. Lawson, The effect of the growth rate on the concentration of nitrogen and transition metal impurities in HPHT synthetic diamonds, Diam. Relat. Mater. 11 (2002) 204. [74] V.D. Blank, V.N.N. Denisov, A.N. Kirichenko, M.S. Kuznetsov, B.N. Mavrin, Raman scattering by defect-induced excitations in boron-doped diamond single crystals, Diam. Relat. Mater. 17 (2008) 1840. [75] H. Sumiya, S. Satoh, High-pressure synthesis of high-purity diamond crystal, Diam. Relat. Mater. 5 (1996) 1359. [76] Y.V. Bataleva, Y.N. Palyanov, Y.M. Borzdov, I.N. Kupriyanov, A.G. Sokol, Synthesis of diamonds with mineral, fluid and melt inclusions, Lithos 265 (2016) 292. [77] H. Kanda, T. Ohsawa, O. Fukunaga, I. Sunagawa, I. Sunagawa (Ed.), Morphology and Growth Unit of Crystals, Terra, Tokyo, 1984. [78] C.J. Widmann, W. Muller-Sebert, N. Lang, C.E. Nebel, Homoepitaxial Growth of Single Crystalline CVD-diamond, Diam. Relat. Mater. 64 (2016) 1. [79] H. Kanda, T. Ohsawa, Effect of solvent metals upon the morphology of synthetic diamonds, J. Cryst. Growth 94 (1989) 115. [80] I. Sunagawa, Growth and morphology of diamond crystals under stable and metastable conditions, J. Cryst. Growth 99 (1990) 1156.

C H A P T E R

2.3 Heteroepitaxial growth of diamond Xin Jiang Institute of Materials Engineering, University of Siegen, Siegen, Germany

Epitaxial film growth is not only of scientific interest but also of high industrial importance. Nowadays, heteroepitaxy is well known as a special technology of semiconductor production used in the context of semiconductor components and microelectronic circuits. After the successful heteroepitaxial nucleation of diamond on 3C-SiC epitaxial films and on bare silicon wafers, respectively, in 1992, considerable technical progress and scientific understanding on the area of heteroepitaxy of the CVD diamond have been achieved. The decisive role of ion bombardment for bias-enhanced epitaxial diamond nucleation is demonstrated. A parameter window for the nucleation time and the kinetic energy of the growing species reflects the importance of a precise process control. Cross-sectional high-resolution electron microscopic (XHREM) investigations have shown detailed interface structure and crystallographic relations between the diamond films and silicon substrates. In addition to the significant improvement of the crystallographic perfection, large-area diamond films of single-crystalline quality have been prepared.

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135

2.3.1 General remarks The last 2 decades have seen the rapid development of metastable diamond growth by means of chemical vapor deposition. Concurrently, a fast-growing interest in diamond technology has been spawned. Combining a wide variety of outstanding material properties [1, 2], diamond can be regarded as one of the most important materials for applications ranging from tribological, optical, thermal, and electrochemical to electronic applications. Taking the electronic applications as examples, diamond-based devices using a hydrogen-terminated surface have been prepared since 2006. Element Six from Europe reported a 1.26-W output power of diamond microwave power devices using a hydrogen surface treatment, with a carrier mobility reaching 150 cm2/Vs and a two-dimensional hole gas concentration of 5 1012 cm2. In 2014, Kawarada et al. [3] reported that a diamond MOSFET device could work at a temperature of 400oC with a breakdown voltage of 500 V and a current density higher than 40 mA/mm2. However, both the hydrogen-terminated surface and the metal-diamond Schottky interface suffer from low carrier mobility and poor interfacial quality with increasing working cycles. For a series of fascinating and technologically relevant applications, large-area and high-quality epitaxial monocrystalline films with controlled doping are still required. With the development of various chemical vapor deposition (CVD) methods during the past 2 decades, an economic preparation of synthetic diamond films has become possible. In general, there are two ways to achieve large-sized single-crystal diamond films. One is the mosaic seeding method, which uses small single-crystal diamond seeds to be properly aligned to form a complete substrate, and then grow diamond homoepitaxially thereon. This method requires that crystal features such as crystal orientation, declination, and height at the adjacent diamond seed contact interface are highly matched. Single-crystal wafers of 40  60 mm2 have been successfully synthesized by some research groups via ion implantation delamination and repeated deposition. It is, however, difficult to achieve such highprecision matching by processing single-crystal diamond seeds, and the grain boundaries of the films prepared are still significant, resulting in deteriorated carrier transport and unstable strength. The other path to achieve large-sized single-crystal diamond films is the heteroepitaxial growth method, which is dependent on the relationship between the atomic arrangements of the substrate and the film. The crystal perfection of the film depends on the physicochemical nature of the substrate surfaces. The successful heteroepitaxial nucleation of diamond on ß-SiC epitaxial films [4] and bare silicon wafers [5, 6] was reported in the early 1990s. Thereafter, considerable technical progress and scientific understanding of the heteroepitaxy of the CVD diamond have been achieved. The decisive role of ion bombardment for the bias-enhanced epitaxial diamond nucleation has been recognized. It has been shown that a narrow parameter window exists for the growth of well-oriented diamond films, especially with respect to the choice of nucleation time and kinetic conditions during deposition. A precise control of the nucleation was found to be necessary for epitaxy. The epitaxial crystal growth is of the Volmer-Weber type, leading to polycrystalline film deposition with epitaxial orientation of individual crystallites with slight misorientation. Even though the special merging of {100} growth sectors of neighboring oriented grains was recognized to terminate the small-angle grain boundaries by a proper introduction of lattice disclination [7], the preparation of large-area single-crystalline diamond films still remains one of the greatest challenges in diamond technology.

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In recent years, film preparation on substrates with a multilayered single-crystal transition layer by a bias-assisted high-density nucleation method has made significant progress in improving the orientation perfection toward the synthesis of large-sized diamond films of single-crystalline quality [8]. The main problems remaining are the controllability of the interlayer stress and nucleation sites as well as the large amount of crystal orientation mismatch that still exists in the prepared single-crystal diamond. The method of reducing the defect density of the diamond by epitaxial lateral overgrowth (ELO) has again drawn people’s attention and masked the defects in the seeds by a metal coating of certain shape and size. Significantly reduced defect density in the epitaxial layer was successful. However, the grain boundary stress shows a continuous deterioration trend due to the introduction of heterogeneous metals. In this session, the advances in heteroepitaxial nucleation and textured growth of diamond will be reviewed and the problems, which remain to be resolved, will be discussed. We will especially focus on the effect of bias-induced ion bombardment on the process of deposition and chemical etching during heteronucleation and selected lateral growth. Recent advances in the deposition of large-area single-crystalline diamond films will be debated. Furthermore, the nucleation process on Ir interfacial layers will be compared to that on Si.

2.3.2 Special problems of heteroepitaxial diamond nucleation 2.3.2.1 Practical significance of nucleation The deposition process of a solid film is divided into nucleation and growth stages. There are three different growth modes for all types of films including island (Volmer-Weber), island- plus- layer, and layer- by- layer growth, which depend on the sum of Gibbs energy and the relative surface energy of the film according to thermodynamics. Diamond films usually feature a Volmer-Weber growth mode due to their high surface energy [5.3–9.2 J/m2, Field] relative to a foreign substrate [2]. This high surface energy is also responsible for a high nucleation barrier for diamond crystals. A low nucleation density (<106 cm2) is found in deposition experiments without any specific pretreatment. Nucleation is the first and critical step of the deposition of heteroepitaxial diamond film. For the deposition of a continuous diamond film, nucleation densities of 109–1011 cm2 are necessary. Moreover, oriented diamond crystals in high nucleation density will facilitate the growth of single-crystalline diamond films.

2.3.2.2 Methods for diamond nucleation A large variety of surface pretreatment strategies have been developed to enhance the nucleation of diamond. Initially, diamond single crystals were used as substrates [9–11]. Later, Matsumoto [12] and Mitsuda et al. [13] achieved a breakthrough in the enhancement of diamond nucleation on nondiamond substrates using diamond seeds and substrate scratching with diamond powder, respectively. Thereafter, substrate scratching has become the most common and powerful method for achieving a high nucleation density and fine grains of uniform size. For silicon substrates, which have been studied intensively, a nucleation density up

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137

to 1011 cm2 can be routinely obtained with substrate scratching while the nucleation density only reaches 106 cm2 without pretreatment. However, it is difficult to achieve oriented heterogeneous nucleation or epitaxial growth on nondiamond substrates through this approach. In 1991, Yugo et al. reported the bias-enhanced nucleation (BEN) method. With the application of a negative bias potential, a high density of nucleation (109–1010 cm2) was obtained on a mirror-polished Si substrate (without scratching) using the MWCVD system [14]. Subsequent developments of BEN by Jiang et al. and Stoner et al. have led to the heteroepitaxial growth of diamond on silicon [5, 6, 15, 16] and on silicon carbide substrates [4, 17], respectively. For the HFVCD method, an enhancement of diamond nucleation similar to that of MWCVD could be achieved under the proper negative substrate bias [18, 19]. BEN treatment has been proven most successful for heteroepitaxial diamond deposition. In the following, the nucleation schemes for deposition on Si and on ß-SiC will be discussed.

2.3.2.3 Heterogeneous nucleation of oriented diamond BEN treatment was first applied to a mirror-polished Si substrate in an MWCVD system. Prior to BEN treatment, in situ hydrogen plasma etching was performed in order to remove the native surface oxide layer. For the nucleation process, the crucial parameters are the substrate temperature, the methane concentration, the applied bias voltage during nucleation, and the nucleation time. A critical bias voltage exists beyond which the energy barrier for the formation of stable nuclei can be overcome [15, 16]. The experiments on Si substrate typically showed a critical bias voltage of approximately 80 to 100 V, depending on other parameters, above all the methane-in-hydrogen concentration and the total process gas pressure. For the effect of nucleation time, investigations by atomic force microscopy (AFM) and reflection-high-energy-electron diffraction (RHEED) revealed that the diamond nuclei formed at the start of the nucleation are heteroepitaxially oriented [20, 21]. As the nucleation time increases, the nucleation density increases dramatically to 5  1010 cm2 and the nuclei become randomly oriented. These results imply the importance of controlling the nucleation process during epitaxial film growth [22]. The investigations mentioned above indicate a narrow parameter window for heteroepitaxial nucleation. Schreck et al. studied this processing window for nucleation of diamond on Si (001) using the BEN process by X-ray diffraction texture measurements [23]. Fig. 2.3.1 shows schematic diagrams of the process parameters. The temporal development of the azimuth pole density distribution in Fig. 2.3.1A demonstrates that the bias time lies within a distinct time interval. As the bias voltage j Vb j decreases, the biasing time and the time window for an optimal crystal alignment increase sharply (Fig. 2.3.1B). By carefully controlling the nucleation process, diamond nucleation with more than 90% [001]-oriented nuclei could be achieved [22]. A possible mechanism of the loss of diamond epitaxy beyond the parameter window for optimized diamond growth is suggested by Jiang et al. and Schreck et al. based on investigating diamond growth under bias conditions [21, 23]. The authors found that even for homoepitaxial diamond growth that shows a perfect lattice match, a misorientation of the subsequently grown crystallites occurs once the underlying crystal lattice is damaged. This means that the orientation deviation of diamond crystals can be traced back to either the formation of defects during the homoepitaxial growth of the crystallites induced by ion

Isatt~50mA tsatt~15min

25

DI= 5.6mA 4mA 2.5mA 1mA

20

15 0

2

4

6

Time (min)

8

10

6min

3

Intensity (arb. units)

FIG. 2.3.1 (A) Azimuthal {220} pole X-ray intensity for diamond films obtained for a variation of the biasing time; (B) Variation of the process time window for oriented nucleation with the bias voltage [23].

Current (mA)

2. Semiconductor diamond

12 FWHM =

2min

12min

8.5°

11.5min

7.1°

11min

6.7°

10min

4.6°

7min

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5min

5.7°

2

1

0 –15

–10

(A)

–5 0 5 Azimuthal angle (°)

10

15

2000 1800 1600 1400

Time (s)

138

1200 topt

1000 800 600 400 200 0 –300

(B)

–280

–260

–240

–220

BIAS VOLTAGE (V)

–200

–180

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139

bombardment or the renucleation of strongly misoriented grains. The lattice distortioninduced misorientation of heteroepitaxial diamond films on silicon and other kinds of substrates represents the reason for the narrow parameter window for a successful biasenhanced epitaxial nucleation.

2.3.2.4 Mechanisms of the bias enhanced nucleation The most-accepted mechanism for bias-enhanced epitaxial nucleation is based on a shallow ion implantation model [24–26]. In this model, the sp3 bonded carbon clusters, formed by low energy ion implantation, act as nucleation precursors. The positively charged ions in the growth chamber accelerated by the negative bias bombard the substrate surface, remove the contamination, and facilitate cluster formation at the surface, which in turn advances diamond nucleation. Stoner et al. reported that the change in plasma chemistry, such as the increase in the concentration of atomic hydrogen caused by substrate biasing and the formation of a carbide surface layer, play an important role for nucleation [27]. Jiang et al. [21] found that the temporal evolution of the overall nucleation density corresponds well with a surface kinetic model established by Tomellini et al. involving immobile active nucleation sites, germs, and nuclei [28]. They also revealed that the enhanced surface diffusion and sticking probability of carbon on silicon due to ion bombardment should be the decisive factors as well as surface defects (point defects, steps, and sp3 bonded carbon clusters) serving as the nucleation sites. In the following, the effect of substrate biasing on diamond nucleation will be summarized based on the role of ion bombardment and surface diffusion. 2.3.2.4.1 Role of ion bombardment for the bias-enhanced diamond nucleation To identify the effect of substrate biasing, BEN treatment was carried out on topographically structured silicon wafers containing grooves having vertically directed side walls prepared using reactive ion etching (Fig. 2.3.2) [29, 30] as well as SiO2-masked silicon substrates [31, 32]. The height of the vertical walls was arranged to be only about 2 μm. With a negative electrical potential applied to the substrate, an ion flux with a preferred direction perpendicular to the silicon wafer will be formed under the acceleration of the positively charged ions in the plasma. The ions flying toward the wafer can’t “feel” the grooved surface structure because of the mean free path of the ions of approximately 5 μm (larger than the structure of 2 μm) under gas pressure around 25 mbar. It is therefore expected that the faces that were parallel to the wafer surface will be bombarded while the side walls hardly experienced any ion bombardment. As a result, diamond nuclei were formed on the surfaces parallel to the wafer and the vertical side walls remained clean. This confirms that the directional flux of energetic species toward the wafer surface is essential for the observed nucleation phenomena. 2.3.2.4.2 Kinetics during bias-enhanced nucleation Jiang et al. demonstrated that the temporal evolution of overall nucleation density agreed with a surface kinetic model involving immobile “active sites,” “germs,” and “nuclei,” [33] which was proposed by Tomellini et al. to describe the formation of diamond crystallites on Si [28].

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FIG. 2.3.2

Selective growth of diamond crystals on a grooved silicon wafer containing both faces perpendicular and parallel to the substrate. The different nucleation densities on the faces confirm the decisive role of ion bombardment for the diamond nucleation. Lower part: sketch of substrate surface and incident ions [29].

The nucleation density N (islands/cm2) is plotted in Fig. 2.3.3 as a function of nucleation time t, which is counted from the induction time tind, a time that the aggregates formed initially have been detectable. The induction tind can also be understood as the period for production of the nucleation sites. It is evident that once the nucleation begins, the nucleation density increases rapidly to 4.3  1010 cm2 within 8.5 min. With increasing time and the

FIG. 2.3.3 Nucleation density N(t) (islands/cm2) versus deposition time t. In the abscissa, an induction time of 6.5 min is subtracted. The curves are obtained by computer modeling using the three-step kinetic models proposed by Tomellini and Polini. The open and full circles in the plots represent data calculated from crystal size distribution (inset) and are obtained by direct particle counting, respectively [33].

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141

increase of density, the main part of the substrate surface is covered by the nuclei. The nucleation rate is then lowered with increasing coverage. From Fig. 2.3.3 it can be noticed that the nucleation kinetics derived from the size distribution function (open circles in the plots, deposition time ¼ 12 min) are close to the result directly measured by particle counting at different times (full circles) and the diamond nucleation agrees very well with the kinetics of the established surface process. 2.3.2.4.3 Surface diffusion during nucleation: measurements of first-nearest-neighbor distance distribution of the nuclei Evidence of surface diffusion of growth species is obtained by measurements of firstnearest-neighbor distance distribution of the nuclei (Fig. 2.3.4, [21]). Considering that surface diffusion of mobile species (adatoms) toward immobile nuclei [acting as sinks for the mobile species, Fig. 2.3.4E] could reduce the formation probability of new nuclei around the immediate vicinity of those immobile ones, the first-nearest-neighbor distance distribution of the nuclei should therefore deviate from a random distribution. Actually, for each nucleus, a depletion zone that corresponds to the region from which a nucleus is not possible to be formed (influenced by the existing ones) is observed in SEM and AFM observations. It implies that the diamond nucleation is strongly influenced by the existing nuclei—an effective “repulsive interaction” among the clusters exists and the clusters tend to “separate.” The surface diffusion of the adsorbed species leads to a decrease of the precursor concentration (Fig. 2.3.4E) and the nucleation frequency drops down to zero on the border of an established island. Fig. 2.3.4A–D compare the measured distributions with the calculated distributions (solid lines) of the first nearest neighbors of a random distribution of diamond nuclei deposited under 150 V, 180, and 250 V, respectively. The nucleation densities of these samples were measured to be 1.1 1010 cm2, 0.68  1010 cm2, and 0.65 1010 cm2, respectively. A depletion zone around each cluster was observed in all the measured distributions. The depletion zone of the nuclei increases from 12 nm for Vb ¼150 V to 35 nm for Vb ¼250 V [33]. Due to ion bombardment under higher energy, accelerated by bias voltage, the mobility and the diffusion of the adsorbed species on the substrate surface will increase, which leads to the increase of the spatial extension of the surface depletion zone. This result confirms the contribution of surface diffusion to the depletion. According to the above analysis, the nucleation sequence by BEN on Si substrate can be summarized as the following: (i) Formation of nucleation sites. By applying a substrate bias, an increased ion bombardment results in a shallow implantation of C if the ion energy is large enough (>20 eV) [25]. A growing amount of C and Si interstitials leads to densification and localized stress, which increase the local surface energy and nucleation. Moreover, an increase of the sticking coefficient of hydrocarbon radicals on the substrate surface promotes the formation of high energy nucleation sites. (ii) Formation of carbon clusters due to the enhanced surface diffusion under a negative bias potential. The enlargement of the first-nearestneighbor distance provides strong evidence for the adatom diffusion. (iii) Formation of stable diamond nuclei. Under a negative bias voltage, the attraction of the positively charged hydrocarbon species leads to a carbon supersaturation at the substrate surface. The growth stage begins once the nuclei reach a critical size.

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Concentration of adsorbed species

Nucleus

(E)

FIG. 2.3.4

Comparisons of measured nearest-neighbor-distance distributions and the distributions predicted from a random nucleation model for samples prepared under different substrate bias voltages. (A) Vb ¼150 V; (B) Vb ¼180 V; (C) Vb ¼250 V; (D) depletion distance versus Vb; (E) schematic illustration of the diamond nucleus as a sink for adatoms [21, 33].

143

2.3.3 The role of substrate material

2.3.3 The role of substrate material Heteroepitaxial growth of diamond films requires appropriate foreign substrates. The selection of substrate materials is mainly based on the principle of minimum system energy and its lattice match to that of diamond. Those substrates featuring relatively high surface energy and matching lattice type and a parameter with diamond (face-centered cubic fcc lattice with ˚ ) will lead to an epitaxial nucleation of diamond crystals. In lattice constant aD ¼ 3.567 A addition, the match of thermal expansion coefficient between the film and substrate will decrease the residual stress and crystalline defects in the films. As a comparison, Table 2.3.1 illustrates these parameters of the most popular substrate materials. All the mentioned materials have the same face-centered lattice as diamond, which is shown by a unit cell in Fig. 2.3.5.

2.3.3.1 c-BN c-BN is the most suitable substrate for diamond heteroepitaxy because of its good match of ˚ ; lattice misfit to diamond ¼ 1.3%) and its lattice parameter with diamond (ac-BN ¼ 3.6120 A TABLE 2.3.1

Parameters of the most popular substrate materials

Material

Lattice constant

Surface energy

Diamond

˚ 3.5667 A

ca. 6.0 J/m2

c-BN

˚ 3.6120 A

4.8 J/m2

ß-SiC

˚ 4.3590 A

2.9 J/m2

Silicon

˚ 5.4388 A

1.5 J/m2

Ni

˚ 3.5240 A

2.4 J/m2

Pt

˚ 3.9230 A

1.4 J/m2

Ir

˚ 3.8390 A

1.8 J/m2

FIG. 2.3.5

aD = 3.567 Å

A face-centered cubic lattice of a diamond crystal.

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relatively high surface energy [γ (111)c-BN  4.8 J/m2; De Vries [34]]. A high nucleation density of diamond crystals of about 1010 cm2 can be achieved without any surface pretreatment of c-BN and without biasing the substrate. Indeed, the first heteroepitaxial diamond growth was reproducibly demonstrated on c-BN (Jones and Gunnarsson [35]; Koizumi et al. [36]; Koizumi and Inuzuka [37]; Wang and Angus et al. [38]; Tomizuka and Shikata [39]; Argoitia et al. [40]). The nucleation and growth of CVD diamond layers show different behaviors on N- and Bterminated {111} faces, caused by the different B-C and N-C bonding strengths. While boron-terminated faces show high diamond nucleation density, the diamond crystals are hardly formed on the nitrogen-terminated faces [37]. The diamond crystals grown on c-BN faces show a parallel, cubic-to-cubic orientation relationship. Unfortunately, a large c-BN single crystal cannot be synthesized at present. Cubic BN films have been grown in a polycrystalline form with an extremely small grain size (several nanometers) by PVD processes. Such films are, however, useless as substrates for the diamond epitaxy.

2.3.3.2 ß-SiC The second substrate candidate chosen for diamond heteroepitaxy is β-SiC. One of the most obvious advantages of ß-SiC in comparison to c-BN is the availability of large single crystals. A large-area epitaxial deposition of ß-SiC onto a silicon single-crystal wafer is also possible. Stoner and Glass obtained highly oriented diamond crystallites using the BEN method, in ˚ ; see Table 2.3.1). Kawarada spite of the large lattice misfit of about 22% (aß-SiC ¼ 4.359 A et al. further improved the nucleation and growth process and obtained continuous [001]oriented diamond films with a high orientation perfection [41]. The X-ray pole figure of the film also shows a cubic-to-cubic orientation between the diamond and ß-SiC substrate, with the full width at half maximum (FWHM) value of the intensity peak being smaller than 1 degree.

2.3.3.3 Si Silicon, as the basic material for today’s microelectronics, is a promising candidate as a substrate material for thin-film diamond devices. Due to the large mismatch of lattice parameter (52%) between diamond and silicon and the much lower surface energy of silicon (γ (111)Si  1.5 J/m2), the heteroepitaxial growth of diamond on Si had been difficult to achieve until 1990. Only randomly oriented island growth of diamond was observed in most of the experiments (Narayan et al. [42]; Jeng et al. [43]). In 1992, Jiang and Klages reported that [001]oriented diamond films could be epitaxially grown on the (001) silicon substrate in an MWCVD process by applying a negative electrical potential to the substrate, as shown in the SEM image in Fig. 2.3.6 [44–46]. They revealed that the interfacial layer of ß-SiC is not necessary for the heteroepitaxial growth of diamond on Si. Later, the structural quality of films on (001) Si was gradually improved with the orientation from a best value of about 9 degrees in 1992 [6] to a best value of about 2 degrees in 1998 [47].

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145

FIG. 2.3.6 Scanning electron microscopic image of a [001]-oriented, 12-μm-thick diamond film on Si [44].

2.3.3.4 Metallic substrates The deposition of diamond films on metallic substrates with a close lattice match has been performed since the beginning of CVD diamond technology. The epitaxial nucleation of diamond on Cu by ion implantation failed [48]. The nucleation of [111]-oriented diamond films could be achieved on Ni and Pt by seeding diamond particles and subsequent plasma annealing [49, 50]. Later, superior-quality heteroepitaxial growth of diamond on oxide substrates with iridium (Ir) transition layers was reported (Ohtsuka et al. [51]; Ohtsuka et al. [52]; Schreck et al. [53]; Schreck et al. [54]). Because of its close lattice match with diamond ˚ ) and its resistance to the formation of a carbide interfacial layer, Ir should be (aIr ¼ 3.840 A a good candidate as a substrate for diamond heteroepitaxy. In 2001, Schreck et al. reported that the spread for the crystal orientations was less than 1 degree [55]. They later realized the heteroepitaxial growth of diamond films with a diameter of 92 mm on Si substrates with Ir as an interfacial layer. In summary, as described in the above sections, the deposition of (001)-oriented diamond films with minimum misorientations of 2 degrees and 1 degree is achieved directly on silicon and ß-SiC, respectively, although there are large mismatches between the diamond films and the substrates. With a long-term effort in optimization of the film growth process, large-area diamond films with strongly reduced defect density have been realized on the Ir/YSZ/Si substrates. Further understanding of the interfaces and defect controls in heteroepitaxial diamond films on Si, ß-SiC, and/or Ir/YSZ/Si substrates will advance the deposition of large-area epitaxial diamond films. Therefore, in the following sections, the interface and defect control of diamond films directly grown on Si substrates will be discussed. After that, the recent progress in nucleation and growth of diamond films on Ir/YSZ/Si substrates will be introduced.

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2.3.4 Diamond-silicon (001) interfacial issues 2.3.4.1 Cross-sectional TEM investigation HRTEM investigation of the interface of a heteroepitaxial diamond film on Si reveals the cubic-to-cubic orientation relationship, that is, (001)diamond//(001)silicon and [110]diamond// [110]silicon as shown in Fig. 2.3.7 (Jiang and Jia [45]; Jia et al. [46]; Jiang and Jia [56]). Obviously, the large lattice mismatch between silicon and diamond (52%) is accommodated by the introduction of a 3:2 registry for the diamond lattice with respect to that of silicon, that is, every three Si (111) fringes are matched well with four diamond (111) fringes, which dramatically reduces the multiple mismatch to about 1.5%. Every two 60-degree interface dislocations meet each other, forming a so-called Lomer dislocation (see Fig. 2.3.7). Diamond crystals are observed to be epitaxially grown directly on silicon, with no formation of secondary phases such as ß-SiC, graphite, and amorphous carbon at the interface area. This observation clearly answered the question of whether the formation of a silicon carbide transition layer was necessary for diamond growth on silicon. Numerous experiments on diamond nucleation confirm that diamond can nucleate directly on an Si substrate if the BEN process is performed at moderate bias voltage and growth temperature while the incubation time is short [57]. This phenomenon, however, is different from the early observation that diamond is grown on Si through means of a β-SiC intermediate layer [58]. After considerable effort, a cross-section HRTEM lattice observation close to a nucleation site was achieved [56]. Fig. 2.3.8A shows the corresponding diamond film with a hillock of a lateral size of only 10 nm on the silicon substrate, as denoted by the white arrow. The morphology of the hillock with facets is traced by the dotted line in Fig. 2.3.8B. {111} twinning occurs on the top layer while the majority of the lamellae start at the facet edges on the sides

FIG. 2.3.7 A [110] lattice-fringe image of the diamond-silicon interface. The multiple lattice fit and the 60 degree and Lomer dislocations are also shown.

2.3.4 Diamond-silicon (001) interfacial issues

147

FIG. 2.3.8 (A) A low magnification image of a diamond grain. An arrow shows a hillock at the central part of the interface. The majority of twins occur near the hillock. (B) Enlargement of the hillock area in (A). A dotted line traces the facet morphology of the hillock. Three arrows mark the facet edges on the sides of the hillock, where the {111} lamellae start [56].

of the hillock denoted by arrows and run into the grain. It is obvious that a direct bonding of the diamond lattice with the silicon substrate is evident at the top of the hillock. This implies that the top of this hillock can be regarded as the nucleation site of diamond. A clockwise tilt of the diamond lattice by 1 degree can be calculated by measuring the {111} fringes with respect to the silicon substrate.

2.3.4.2 Experimental investigations of crystal misorientation As mentioned in the previous sections, the crucial issue preventing single-crystal diamond growth is the formation of small-angle grain boundaries induced by the orientation deviation of individual diamond grains with respect to the silicon substrate. Measurements by X-ray rocking curves demonstrate a slight orientation deviation with line widths of up to several degrees in the [100]-oriented diamond. Corresponding AFM results reveal that the measured crystal tilting of the (001) facets around the [110] crystal axis excellently meets the statistical Gaussian distribution [59]. The reason for this tilt could be explained by the lattice mismatch between diamond and silicon [60]. According to this model, all the grown crystallites will tilt at a certain angle to accommodate the lattice matching. It was observed that the (111) plane deviates for a few degrees from the ideal orientation by a rotation around the common [110]-zone axis. This small deviation is compensated for by varying the number of the two types of {111} planes that terminate at the interface. The number of the terminating planes along (111) and (111) (see Fig. 2.3.7) is varied while their total number is almost unchanged with respect to the case of the ideal orientation. However, the total number of the 60-degree interface dislocations is found to remain constant for a misorientation up to 14 degrees. It is, however, difficult for this model to interpret the statistical tilt distribution and the location of the maximum of the tilting angle distribution at zero degrees. Furthermore, the observation of bias-enhanced CVD diamond films indicates that ion bombardment could lead to a misorientation even for the case of a homoepitaxial growth of diamond where the lattice mismatch is zero [7]. It was found that the surface roughness

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FIG. 2.3.9

High-resolution lattice image of a diamond grain with a tilt angle of about 2 degrees around the [110] axis, showing a small (001) facet on top of the hillock. An arrow denotes a facet edge inducing a twin boundary [56].

of the silicon substrate is one of the most important factors influencing the orientation of the crystalline nuclei. A certain extent of the roughness of a silicon surface such as surface steps is clearly seen in nearly all observed locations. Fig. 2.3.9 shows a HRTEM image of a diamond grain directly heteroepitaxially grown on Si. The top (001) facet of the hillock in the central region of the interface features a dimension of only several (110) atomic planes. The diamond grain in this image shows a tilt deviation of 2 degree from the ideal orientation by a rotation around the common (silicon and diamond) [110] zone axis. Note that a twin boundary is seen once again to start at a facet edge. A larger misorientation is obtained for the diamond grain in Fig. 2.3.9 compared to that seen in Fig. 2.3.8B. The deteriorated orientation of the diamond grain can be attributed to the relatively rough morphology of the hillock on the substrate.

2.3.5 Particular structures and defect control of [001] diamond films

149

Based on the results presented above, it is evident that the misorientation of the diamond grains on (001) silicon is closely related to the appearance of the hillocks on the substrate where the diamond grains directly contact the silicon.

2.3.4.3 Modeling of the causes of the crystal misorientation The above observations provide us with strong evidence that crystal tilting is related to structural distortion of the substrate surface and thus to the concomitant local strain. The tilting induced by the roughening of the substrate surface is obviously very complicated with a range of possibilities. For heteroepitaxial diamond growth on Si (001), however, the local strain is not only caused by surface roughening but also by lattice mismatch. The local strain induced by lattice mismatch may result in crystal tilting, especially when ion bombardment causes surface etching. In order to understand the origin and nature of the experimentally observed diamond grain tilting, the three-dimensional nucleation and growth process were studied in a stepby-step manner by means of the molecular orbital PM3 calculation (Zhang et al. [61]; Jiang et al. [62]), which is based on the MNDO semiempirical Hamiltonian of the Hartree-Fock theory. A cluster model composed of more than 100 silicon atoms with a hydrogen-saturated boundary was selected to simulate a rough silicon (001) substrate involving a silicon island on the surface. The formation of a diamond embryo actually partially compensates the bond tilt created during the deposition of the first carbon layer. Due to the limited contact area between the diamond embryo and the Si (001) terrace, on the other hand, the structure mismatch will be accommodated during crystal growth by the formation of interfacial dislocations at the interface (pentagon/heptagon bonding configurations) and a crystal tilt will remain. Based on this model, the tilt in the HRTEM results of Fig. 2.3.9 can be well understood. However, if the pentagon/heptagon unit is positioned symmetrically in the diamond embryo, the local lattice strain will be balanced in a symmetrical way. This finding implies that if such a diamond embryo is formed in a sufficiently large area of a (001) terrace, the initial tilt should be significantly reduced or eliminated by a homogeneous distribution of the pentagon/heptagon unit. The fact that a relatively flat top of the hillock in Fig. 2.3.8 leads to a better orientation (θ ¼ 1 degree) of the diamond grain corroborates the theoretical model. In practice, the deposition conditions in the nucleation stage should be chosen in such a way that the nucleation rate is as high as possible and the etching effect of the plasma is as small as possible in order to obtain a well-oriented diamond film on a silicon substrate.

2.3.5 Particular structures and defect control of [001] diamond films 2.3.5.1 Microstructure of [001] diamond films Epitaxially [001]-oriented films synthesized by a bias-enhanced nucleation process consist of columnar grains with a lateral size of several hundreds of nanometers to several micrometers (John et al. [63]; Hessmer et al. [64]; Maeda et al. [65]; Tachibana et al. [66]). Fig. 2.3.10 shows a typical cross-sectional SEM morphology of an 8-μm thick [001]-oriented diamond film. It is obvious that the individual grains grow together, leading to the formation of

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FIG. 2.3.10

SEM micrograph of a [001] diamond film showing the coalesced film surface and the columnar film structure.

FIG. 2.3.11 Plane view of an enlarged area of a small-angle grain boundary. Terminating planes/dislocations are indicated by arrows [68].

a columnar structure. By TEM investigations of similarly grown oriented CVD diamond films, the films can be divided into two layers: one with partial randomly oriented grains formed in the near-interface region and another one featuring orientated columnar grains at a film thickness of above a few hundreds of nanometers. This finding can directly be related to the so-called “evolutionary selection” model and potentially explains the reason for columnar growth under the appropriate growth parameters [67]. The fastest growth rate along the [001]-direction leads to grain boundaries nearly parallel to each other and perpendicular to the silicon substrate. The residual stresses induced by the mismatch in lattice constants and thermal expansion coefficients result in the formation of grain boundaries with different misorientation angles. The small-angle grain boundaries can be described by a dislocation model [68]. Fig. 2.3.11 shows such a grain boundary section, taken at a [100]-zone axis orientation. Both grains are ˚ are seen in the [001]-zone axis orientation, and the {220}-lattice fringes with distances of 1.26 A clearly visible. The misorientation between the two grains is about 3–4 degrees and is compensated by the introduction of dislocations (arrows). The average distance between the dislocations is 1.9 nm. Similar to the observations reported for [110]-zone axes lattice images, no inclusions or additional second phases can be observed at these grain boundaries. The small-angle misorientation between grains is observed to be connected with arrays of dislocations in all cases.

2.3.5 Particular structures and defect control of [001] diamond films

151

2.3.5.2 Crystal coalescence during deposition From the cross-section SEM micrograph in Fig. 2.3.12A, the interesting aspect that the lateral overgrowth of diamond grains is achieved by the disappearance of some grain boundaries in the vicinity of a film surface can be derived. Two other growth modes were proposed FIG. 2.3.12 (A) HREM image of a grain boundary showing the disappearance of a 2-degree small-angle grain boundary at B; (B) schematic representation of stopping a low-angle grain boundary by introduction of a disclination [7].

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FIG. 2.3.13 Schematic representations of two [001] grain growth phenomena–coherent coalescence of grains and overgrowth of grains by lateral step flow growth of neighboring grains.

for the enhancement of orientation and lateral grain size [7]. In the first mode (Fig. 2.3.12B), the misorientation of growing grains should first be reduced to such a small value (Van der Drift selection) that the coalescence of grains can occur [69], thus leading to a singlecrystalline top layer. As shown in Fig. 2.3.12A, with respect to the substrate and the grain DI, a 2-degree tilt angle can be measured near the interface for the grain DII. For this 2-degree tilt grain boundary, grains DI and DII are merged into a single crystal at the position indicated by the letter “B” with the disappearance of the grain boundary. This result provides strong evidence that individual [001] diamond grains can coalesce to form a single crystal, in the case that they feature only a slight orientation deviation. In addition, the misorientation for smallangle grain boundaries could be compensated by a disclination with a corresponding rotation angle (Fig. 2.3.12B). In the second mode, a process favoring lateral growth is applied, allowing the grains to overgrow adjacent ones and resulting likewise in crystals with large lateral dimensions. Fig. 2.3.13 summarizes schematically these two growth modes. The second growth mode will be discussed in the following section in more detail.

2.3.5.3 Lateral growth for preparing films with large grains Heteroepitaxial [001]-oriented diamond films with improved lateral grain size and surface smoothness can also be prepared using a [001]-textured growth process followed by a lateral growth step that is based on a variation of the fastest growth direction (Kawarada et al. [41]; Jiang et al. [47]). This variation can be achieved by changing the process parameters, first the substrate temperature as well as the methane concentration. The best lateral growth could be realized by a [110] step-flow growth process of the (001) surface, as demonstrated by the cross-sectional SEM image shown in Fig. 2.3.14 [47]. The results indicate that the diamond crystals increase their lateral dimensions at the (001) film surface by changing the grain boundary plane orientations from preferentially vertical to preferentially parallel directions with respect to the (001) growth faces, as schematically shown in Fig. 2.3.13. The grains with a relatively large angle deviation from the ideal epitaxial orientation are overgrown by those featuring a relatively small angle deviation. As a result, the degree of orientation perfection of the films improves in comparison to that of films prepared by the established process of [001]-textured growth. The presence of boron in the gas phase was found to encourage the step-flow lateral growth. It was possible to achieve the deposition of a thin boron-doped diamond film on silicon, characterized by a full width at half maximum value of the measured tilt angle distribution of only 2.1 degrees.

2.3.5 Particular structures and defect control of [001] diamond films

153

FIG. 2.3.14 Cross-sectional SEM micrograph showing the change of the grain boundary orientations in the boron-doped top layer. The grain boundaries indicated by arrows are overgrown by a neighboring grain; the grain boundary between the two grains is switched.

Step-flow-induced lateral growth can occur if the growing face of the film is completely covered by (001) facets. Such step-flow-induced growth has been observed previously on the (001) face of a single-crystal diamond substrate during homoepitaxy (Lee and Badzian [70]; Hayashi et al. [71]). It was found that multiatomic steps belonging to terraces of approximately 8 nm in lateral dimensions that propagate along the [110] direction of the surface result in net film growth in the perpendicular [001]-direction. SEM and TEM observations [47] indicate that the grain boundary plane orientation is related to the existence of the {111} facets leading to the formation of pits on the growth surface. If the (001) growth faces are separated by the {111} faces, surface pits are formed. In this case, the step flow cannot occur and the grain boundary propagates in a direction vertical to the substrate surfaces. The step flow stops if a step is encountered on another grain that moves in a reverse direction.

2.3.5.4 Structural imperfections and defect control in diamond films Defects at all scales play a key role in material properties. Defect control and/or engineering is also a fundamental issue with respect to the intrinsic properties of the CVD diamond. Structural imperfections of CVD-grown diamond crystals include grain boundaries, stacking faults, dislocations, and interfaces to substrates. Point defects and impurities will be discussed in other chapters. The defects discussed here are growth-related imperfections. TEM cross-sectional investigations show a highly defective film structure in the interface region composed of small domains of multiple twinning and amorphous carbon phases. In a later CVD deposition stage, the (001) textured film growth can be realized by an evolutionary competition process (the so-called van de Drift mechanism) that reduces the film misorientation [72]. Diamond crystallites grow larger and are mainly bounded with {100} and {111} facets, as schematically shown in Fig. 2.3.15A. The {111} growth sectors of [001] films produce a higher density of planar stacking faults and microtwins while no twins and only a few dislocations are formed in {100} growth sectors. These phenomena were observed in cross-section and plan-view micrographs of [001] diamond films after a lateral growth stage. Fig. 2.3.15B is such a cross-section

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(A)

(B)

(C)

FIG. 2.3.15 (A) Schematic cross-section of [001]-oriented columnar grains showing the {001} and {111} growth sectors. (B) Cross-sectional HRTEM micrograph recorded in the [110] direction showing the change of the grain boundary orientation. (C) Coalescence of boron-doped rectangular-shaped grains on the surface of the diamond film. In (B) and (C), the higher density of defects around the grain boundary is clearly observable.

2.3.6 Growth of large-area (001) single-crystalline diamond films

155

TEM image that shows two [001]-oriented diamond grains after a boron doping-induced lateral growth. In a region close to the grain boundary, a higher defect density is observed. The plane-view image of the same sample also shows defects close to the grain boundaries (Fig. 2.3.15C).

2.3.6 Growth of large-area (001) single-crystalline diamond films (001) single-crystalline diamond films of a large size could be obtained when the grain boundaries of a (001)-textured diamond film are removed during the overgrowth process. To do so, a strongly improved and narrower nucleation and growth parameter window is required. As shown in the sections above, the deposition of (001)-textured diamond films on an Si substrate with a FWHM of grain tilt angles of 2.1 degrees could be achieved by tuning the growth parameter to optimize the lateral growth process. However, a high density of small-angle grain boundaries in the (001) plane of these films still remains. These small-angle grain boundaries are induced by the residual stress during the nucleation and growth stage. For the growth on an Si substrate, the mismatch of the lattice constant plays an important role on the distribution of the residual stress. In order to decrease the defect density and the residual stress, the introduction of appropriate interfacial layers on Si as well as the selection of other foreign substrates have been investigated in recent years. Iridium (Ir) is a promising substrate material for the deposition of diamond because it features excellent properties such as a close lattice constant with diamond and a low solubility of carbon. High-quality epitaxial diamond films have been deposited on (001) MgO or SrTiO3 substrates with an interfacial layer of (001) Ir by the BEN method [55]. Also, it is confirmed that diamond films, which have been heteroepitaxially grown on an Si/YSZ/Ir substrate by BEN (i.e., Si substrate with the interfacial layer of yttria-stablized zirconia (YSZ)/Ir), show a higher crystalline quality along with a decreased density of crystal defects compared to those that have been deposited directly on a pure Si substrate [8]. The critical issue in depositing high-quality diamond on Ir is the diamond nucleation stage, which differs from that on Si in some aspects. In the following part, the nucleation process and the defect control of diamond films on Ir will be discussed in more detail.

2.3.6.1 Nucleation mechanism on Ir Under BEN treatment, the heteroepitaxial nucleation of diamond on Ir features distinct differences compared to the corresponding process on an Si substrate. First, diamond nucleation occurs under a slightly higher negative bias voltage toward Ir (such as 250 V instead of 200 V) while the current is typically more stable over time. As mentioned in the section above, the selective etching induced by the bombardment of H+ ions plays a dominant role in the oriented growth of diamond crystals. Applying a higher bias voltage toward the substrate, the etching rate of diamond by H+ ions tends to increase. It is observed that diamond crystals with an initial size of up to 0.5 μm on an Ir interlayer nearly disappear after an exposition to a bias treatment for an hour [73]. In this case, the nucleated crystals are not stable and tend to shrink under H+ ion etching, resulting in inhibited diamond growth while the

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nucleation probability on a large-scale substrate is enhanced. In addition, the nucleation mechanism under etching conditions can’t be explained by classical nucleation theory, where the clusters with a critical size can grow under a positive supersaturation of the precursor gas. Two phenomena are observed during the initial diamond nucleation stage under the ion bombardment of a BEN treatment. First, several groups reported that the corresponding Ir surfaces become rough due to the effect of ion etching (Hormann et al., [73]; Sawabe et al. [74]). Second, a thin carbon layer is formed on the roughened Ir surface. The thickness of this carbon layer is estimated to be about 1–2 nm. This finding is supported by numerous techniques such as elastic recoil detection analysis (ERD, Bauer et al. [75]), X-ray photoelectron spectroscopy (XPS, Kono et al. [76]), AFM (Gsell et al. [77]), and TEM (Brescia et al. [78]), Schreck et al. [79]). Although this layer is thin, it turns out that its morphology is well observable using in-lens detectors in SEM. The SEM images show islands featuring a bright contrast with well-defined borders distributed inside the carbon layer. The contrast itself is homogeneous within individual islands, which are called characteristic domains or patterns. It was confirmed that diamond is nucleated within these domains by means of dedicated short growth experiments immediately following the BEN treatment. Investigation of the phase structure of this carbon layer will certainly advance the understanding of the corresponding nucleation mechanism. The layer may be composed of amorphous sp2- or sp3-carbon featuring an amorphous or crystalline structure. Nevertheless, due to its small normal dimension, a distinction of these phases proves challenging. Various characterization approaches, including AFM, X-ray photoelectron diffraction (XPD), high-resolution TEM (HRTEM), and others, are employed to study its crystalline structure. Conductive AFM measurements [77] showed a high electrical resistance inside and outside the domains and can therefore be utilized to exclude the presence of sp2 carbon within this layer. In addition, the relatively high elastic modulus of the carbon structure measured by lateral force microscopy indicates the presence of sp3 carbon inside the domains [79]. Furthermore, cross-sectional HRTEM investigations only revealed amorphous structures in the nucleation layer formed by BEN. No crystalline phase was found within this layer. This finding agrees well with the results of reflection high-energy and low-energy electron diffraction (RHEED and RLEED) measurements. Here, the diffraction spot corresponding to the diamond phase was not observed (Gsell et al. [80]; Ohtsuka et al. [51]; Kono et al. [76]). Later, XPD, which records the signal of the photoelectrons scattered forward from neighboring atoms in a short range, was employed to investigate the crystalline structure of the nucleated layer. In Fig. 2.3.16, the diffraction patterns from Ir 4f electrons and C 1s electrons are recorded for two types of samples deposited on Ir/SrTiO3 by BEN [80]. The diffraction pattern from Ir 4f signals in these samples are identical compared to the one from the reference Ir sample. This implies that the microstructure of Ir is not varied during the BEN treatment. In comparison with the diffraction pattern of C 1s from a standard diamond phase, the sample that experienced an unsuccessful BEN treatment doesn’t show any crystalline features (Fig. 2.3.16D) while the sample that was successfully treated by BEN shows perfect crystalline features that can be identified as the cubic structure of the diamond phase (Fig. 2.3.16F). The latter finding provides strong evidence that sp3 carbon is crystallized in the layer nucleated by BEN. Combined with no observation of diamond features in the TEM, RHEED, and LHEED studies, it can be confirmed that the sp3 carbon in the layer is crystallized in a short-range order instead of a long-range order. A high density of defects including crystal misorientation is

2.3.6 Growth of large-area (001) single-crystalline diamond films

(A) lr reference

(B) Diamond reference C 1s

lr 4f

(C)

lr 4f

(D)

157

FIG. 2.3.16 XPD patterns of the Ir 4f and the C 1s core levels taken for (A) the Ir reference and (B) the diamond reference sample as well as for (C), (D) BEN sample A and (E), (F) BEN sample B [80].

C 1s

BEN Sample A

(E)

lr 4f

(F)

C 1s

BEN Sample B

generated in this layer, which leads to the blurring of the diffraction spot in the figure. For a specific sample, it was possible to estimate, based on quantitative statistics from XPD C 1s patterns, that about one-fourth of the carbon atoms in the layer are situated in an ordered crystalline diamond environment. Other approaches have been used to reveal the structure of this BEN-induced layer. The spectra recorded by X-ray absorption near the edge structure (XANES) combined with X-ray photoemission electron microscopy (X-PEEM) show a dip around 303 eV, which is a fingerprint for diamond [81]. Moreover, in the Auger spectra, the carbon signals from the domains are identified as diamond while the regions outside the domains show signals from graphitic structures [79]. This provides the picture that the nucleated layer consists of isolated crystalline diamond clusters corresponding to the domains embedded in an amorphous carbon matrix. The spectra obtained by high-resolution electron energy loss spectroscopy reveal that the peak associated with diamond’s optical phonons is observed in the nucleated layer [82]. In addition, investigations of the initial growth of diamond on this nucleated layer have been conducted by experiments featuring a growth duration of 5 s; these were carried out immediately after the BEN treatment, as shown in Fig. 2.3.17. It was observed that isolated diamond crystallites were heteroepitaxially grown on the Ir substrate displaying a crystallographic cube-to-cube relationship of (001)diamond//(001)Ir and [110]diamond//[110]Ir. This

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FIG. 2.3.17

Cross-sectional HRTEM micrographs of diamond grains formed during 5 s growth after termination of the BEN treatment (40–45 min). The 10 nm covering layer was deposited before TEM sample preparation. Both samples were deposited under virtually identical conditions Tsub ¼730– 750°C 7% CH4/H2, UBias ¼260 V. In (A) isolated grains had formed while in (B) a 257-nm wide single-crystal layer was observed [8].

relationship is identical compared to the one found for the epitaxial deposition of (001)orientated diamond films on a (001) Si substrate in previous studies. However, the density of misfit dislocations in diamond is dramatically decreased in the case of the Ir substrate because it possesses a significantly better match of the lattice constants. These diamond grains feature a height of about 2 nm and an average interspacing of about 15  20 nm, corresponding to a nucleation density of 3  1011 cm2. No amorphous carbon film is observed surrounding the isolated diamond crystals. Consequently, it can be concluded that the amorphous carbon layer formed by the ion bombardment in the nucleated layer is etched under standard growth conditions without a bias potential. In this scenario, the crystalline sp3 carbon clusters in short-range order act as nuclei for the diamond growth. Based on the information provided above, the nucleation process of diamond on Ir by means of BEN could be understood as follows: (i) Formation of an amorphous carbon layer with short-range order sp3 clusters. Bias-induced ion bombardment leads to an increased sticking absorption and surface diffusion of hydrocarbon radicals on the Ir surface. The low solubility of C in Ir gives rise to a rapid saturation, which results in the formation of a carbon layer. This stage represents a big difference compared to diamond growth on an Si substrate. The thickness of this layer is typically 1–2 nm and independent of the actual bias time. (ii) Formation of diamond nuclei. By applying a negative bias potential, the subimplantation of hyperthermal particles on the highly defective crystalline matrix leads to the formation of diamond nuclei or embryos, as proposed by Lifshitz et al. [83]. For the combined process of deposition and ion etching, the growth of diamond nuclei along the depth direction is suppressed while the lateral growth is enhanced. This mechanism resembles the one found on Si substrates. (iii) Formation of stable diamond nuclei. The number of nuclei featuring a size larger than the critical one increases while smaller nuclei are etched by

2.3.7 Summary and conclusions

159

ion bombardment. The latter suggestion is strongly supported by the observation that the highly defective crystalline sp3 carbon matrix between the nuclei is absent after switching off the bias voltage.

2.3.6.2 Defect control and growth of single crystalline diamond on Ir It is reported that the diamond crystals nucleated on Ir are epitaxially oriented with respect to the substrate [51] because the mismatch of lattice constants for both partners (diamond and Ir) is small (about 8%). This is illustrated in Fig. 2.3.17. As a result, here the epitaxial nucleation density of diamond crystals reaches a value of 3 1011 cm2, which is higher than that found on Si and SiC substrates. Although the nucleation density on Si potentially can reach a density as high as 1011 cm2 by BEN assistance, the majority of these nuclei will become more or less “misoriented” due to the ion bombardment associated with the BEN process. Only nuclei that are not exposed to such an ion bombardment are epitaxially orientated. In addition, for a diamond film on Ir that features a thickness of about 600 nm, the mosaic spread of epitaxial grains is measured to be about 1 degree. This is much smaller compared to the one observed for a diamond film of similar thickness on an Si substrate. With increasing film thickness, the mosaic spread of epitaxial diamond films is reduced. In this context, a tilt of about 0.2 degrees was observed for 35-μm thick diamond films on Ir. As mentioned above, diamond crystals will merge into larger grains based on the introduction of disclinations during the overgrowth process. Consequently, the orientation deviation of diamond films grown on Ir substrates is dramatically narrowed. The defect evolution in the diamond films grown on Ir was studied for films featuring different thicknesses. At a thickness of 600 nm, the diamond film features a morphology of individual grains with some grooves among them. Once the film thickness increases to 8 μm, diamond forms a continuous film with small-angle grain boundaries that almost exclusively consist of dislocations. At a film thickness of 34 μm, the surface of the diamond film became flatter and the grain size increased. For such a film, the dislocation density decreased significantly and is estimated to be on the order of 5–10 108 cm2. At this point, the formation of a grain boundary network has evolved into short defect bands consisting of clusters of dislocations. Raman spectra reveal that the FWHM of diamond films featuring a thickness of several hundreds of micrometers is less than 3 cm1. Based on the assumption that the average dislocation density follows a 1/t scaling law (with t being the film thickness), a further decrease in dislocation density could be achieved. Schreck et al. has prepared a single-crystalline diamond film with a diameter of about 90 mm on Ir [8]. The thickness of this film was about 1.6 mm, with a dislocation density estimated to be 4  107 cm2. This density is much lower than that observed for standard natural Type IIa crystals (108–109 cm2).

2.3.7 Summary and conclusions Significant progress in heteroepitaxial diamond growth has been achieved since 1992 following the introduction of the bias-enhanced nucleation process. The synthesis of single-crystalline diamond films and their applications as high-temperature electronic material still remain a

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great challenge for materials scientists. The crucial issue preventing single-crystal diamond growth is the crystallographic orientation deviation of individual diamond grains with respect to the corresponding substrate and the resulting small-angle grain boundaries in the films. The reduction of the average misorientation angle between grains in (001)-oriented CVD diamond films has been experimentally achieved by employing special modes of crystal growth. It was found that diamond grains with a very small misorientation angle of about 2 degrees can coalesce to form larger diamond grains by terminating their small-angle grain boundary based on a lattice disclination. Furthermore, it could be demonstrated that the presence of boron in the gas phase during film deposition favors a step-flow lateral growth mode. This mode leads to an overgrowth of diamond grains featuring a relatively large misorientation by their highly oriented neighbors. In this context, changes of the grain boundary orientation from an inclined, nearly perpendicular direction to a more parallel direction with respect to the substrate surface were observed by TEM. Here, the parallel direction indicates the overgrowth of neighboring diamond grains. Films prepared in such a way show an improved smoothness of their surface and an increased lateral grain size. In order to realize a large-area high-quality epitaxy, further efforts are required. For example, in order to achieve such heteroepitaxy of diamond on silicon, strategies based on an improved surface treatment and an exact control of nucleation must be taken into consideration. Here, surface treatment is the first and probably the most essential requirement. The substrate surface must be clean as well as free from surface contamination and oxidation. Furthermore, the dangling bonds of the Si surface atoms must be saturated by hydrogen. During the process of bias-enhanced nucleation, the surface roughening due to ion etching and the formation of amorphous carbon must be avoided. The bias voltage, biasing time, and pressure strongly influence crystal orientation. The bombardment with ions possessing an energy beyond a certain critical value is necessary for the formation of nuclei. Unfortunately, a negative influence of the ion bombardment on the alignment of the diamond grains can also be demonstrated. Consequently, an improved epitaxy requires a compromise between the positive and negative aspects of the bombardment. Additionally, as the base pressure of the deposition chamber during CVD diamond nucleation and growth is low at present, any residual gas in the chamber would contaminate or even oxidize the Si surface and change its surface status. Therefore, an improved base vacuum in the growth chamber as well as a higher purity of the gas source would also be beneficial with respect to achieving better epitaxy. In contrast to this, the ion bombardment and etching during the BEN process on an Ir-based substrate does not influence its surface states negatively. It also favors the lateral growth of a carbon layer and suppresses vertical diamond growth completely, which results in the formation of a special domain. The lateral growth of the nucleation layers even facilitates the preferred alignment of diamond crystals featuring a low angular spread. The excellent arrangement of epitaxial crystals on Ir during the nucleation stage leads to the reduction of the mosaic spread of diamond crystals in the subsequent growth of a corresponding diamond film.

Acknowledgments The critical reading of and suggestions relating to this chapter by Dr. Bing Yang and Dr. Thorsten Staedler are especially appreciated.

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C H A P T E R

2.4 Homoepitaxial growth of single-crystal diamond Peng Jina,b, Wangcheng Yua,b, Ye Zhanga,b, Ju Wua,b, Zhanguo Wanga,b a

Key Laboratory of Semiconductor Materials Science and Beijing Key Laboratory of Lowdimensional Semiconductor Materials and Devices, Institute of Semiconductors, Chinese Academy bCenter of Materials Science and Optoelectronics Engineering, University of Chinese Academy of Sciences, Beijing, P. R. China

As a semiconductor material, diamond has many superior physical properties such as an ultrawide bandgap (5.5 eV), high thermal conductivity (22 W/cm
K), high dielectric breakdown field (10 MV/cm), and high carrier mobility (4500 cm2/V
s for electrons and 3800 cm2/V
s for holes in room temperature [1]) as well as excellent optical, infrared, and X-ray transparency. These remarkable properties make diamond a promising candidate for a large variety of device applications, such as high-frequency field-effect transistors, diodes and switches to operate at high temperature, radiation detectors, etc. However, the realization of these devices with excellent performance as expected is not an easy task, as high-quality diamond crystals, effective doping in diamond, metal-diamond contacts and other basic problems in material growth and device processing have been difficult to obtain. Although diamond has tremendous potential for development, the reserves of natural diamonds are very scarce. Therefore, the synthesis of man-made diamonds has been a hot topic,

2.4.1 CVD homoepitaxial growth of diamond

165

both in academia and industry. For decades, scientists have been devoted to the synthesis of diamond, not only to reduce the cost the diamond material but also to utilize the excellent properties. Nowadays, the synthetic diamond methods mainly consist of two techniques: high temperature high pressure (HPHT) and chemical vapor deposition (CVD). This chapter is a brief illustration of the homoepitaxial growth of diamond via the CVD method.

2.4.1 CVD homoepitaxial growth of diamond The diamond fabrication techniques can be divided into two categories: the high-pressure, high-temperature (HPHT) synthesis from graphite and epitaxial growth on substrates via chemical vapor deposition (CVD). The HPHT method is based on the phase diagram of carbon, which indicates that diamond is more stable than graphite under high pressure and high temperature (1300–1700°C and 5–6.5 GPa). The first successful HPHT synthesis of diamond was achieved by the GE Company using graphite and nickel under high temperature and pressure conditions in 1955. However, most diamonds produced by the HPHT method are small grains containing some nitrogen impurities (from the air) and metal impurities (from the metal catalyst), which seriously affect the electrical and optical quality of the diamond. Therefore, these can mainly be used for the preparation of tool-grade diamond material used in grinding and other abrasive applications. In the CVD method, in comparison to HPHT synthesis, diamond crystals of both high crystallinity and significantly larger areas can be epitaxially grown on the substrates, heteroepitaxial or homoepitaxial. In addition, due to the absence of solvents/catalysts used in HPHT synthesis, the CVD diamond is relatively more pure. Therefore, CVD is the main method for the fabrication of diamonds potentially to be used in various semiconductor devices. There are a number of CVD variants for diamond growth: microwave-plasma CVD (MPCVD), hot-filament CVD (HFCVD), DC plasma CVD, and RF plasma CVD. Among these different CVD methods, MPCVD is the most popular for the fabrication of diamond crystals as it has the following characteristics: high plasma density, high stability and reproducibility, high purity deposition environment without electrodes, and the potential for large area growth. In the MPCVD growth of diamond, both diamond (i.e., HPHT diamond) and foreign materials (e.g., iridium) can be used as substrates. However, although the heteroepitaxial growth on foreign substrates is a promising approach for large diamond materials, due to the lattice mismatch between the diamond and the substrate, the crystallinity of the heteroepitaxial diamond is still largely restricted by some important problems such as the residual mosaicity, the residual strain, and the presence of a high density of dislocations [2]. Therefore, the CVD epitaxial growth of high-crystal-quality diamond is mainly focused on the homoepitaxial growth in both fundamental investigations and applications. The CVD epitaxial growth of diamond can be classified as homoepitaxial growth and heteroepitaxial growth. The homoepitaxial growth technique is an easier way to form a high-quality epitaxial layer. HPHT diamonds are often used as the substrate for homoepitaxial growth. However, the HPHT diamond usually has smaller areas, which severely limits the increment of epitaxial diamond size.

166

2. Semiconductor diamond

2.4.2 Growth mechanisms of CVD diamond Unlike silicon or germanium, carbon is capable of forming many allotropes such as diamond, graphite, and nanocarbons. During the growth of diamond, other carbon allotropes could codeposit together with diamond, leading to impure mixed phases. To obtain a highquality diamond, atomic hydrogen or oxygen is usually used for the preferential etching of graphite. Therefore, physical-vapor deposition processes are not suitable for the deposition of diamond. The CVD process, on the other hand, involves gas-phase chemical reactions occurring above the substrate surface, allowing the achievement of selective etching of other carbon allotropes. CVD techniques for diamond growth require the activation of gas-phase carbon-containing precursor molecules. This activation differs in different CVD variants; for example, thermal activation is used in HFCVD, a combustion flame is used in oxyacetylene torch growth, and electric discharge is used in DC, RF, or microwave CVD. As mentioned before, MPCVD is currently the most popular CVD method for diamond fabrication. The MPCVD growth of diamond utilizes plasma to activate reactant molecules, a typical process at the substrate temperatures in the range of 700–1100°C, while the pressures vary from a few Torr to several hundred Torr. According to the carbon phase diagram, diamond is relatively unstable with respect to graphite at the typical MPCVD growth conditions. The successful fabrication of diamond by MPCVD relies on both chemistry and kinetics reactions with the help of plasma. A simplified CVD diamond growth process can be summarized in the following steps. First, different reactant gases are thoroughly mixed together before entering the plasma region. Second, the plasma ball is formed directly above and adjacent to the substrate. As the gas species pass through the plasma region, they are ionized and dissociated by the microwave energy into reactive radicals and atoms. Finally, the reactive radicals reach the substrate, continuing to mix and undergo a series of chemical reactions on the substrate surface. With the appropriate surface reactions and gas chemistry conditions, diamond can be formed. As for the reactant gases, the MPCVD process is based on the activation of hydrocarbon, typically methane, in a mixture with hydrogen. Other carbon precursors, including methane, acetylene, ethylene, ethane, and carbon dioxide, have also been tried as carbon sources in diamond growth although they are rarely used now. Besides, oxygen or nitrogen can be added in the feed gas as catalysts, which could change both plasma parameters and surface chemistry. To determine the composition of input gas mixtures in the MPCVD growth of diamond, the "Bachmann triangle diagram," as shown in Fig. 2.4.1, is proposed as a guideline. This triangle diagram shows experimental gas compositions of hydrogen, carbon, and oxygen under various MPCVD methods. It is defined that diamond growth would happen only when the gas composition ratio was inside the growth region within the diagram, independent of deposition system or gas-phase precursors [4]. However, it should be noted that the correct gas composition is a necessary but not sufficient condition for diamond growth. The correct surface temperature, growth surface state, and other growth parameters are also required for successful growth. At the second step of the diamond growth process, the plasma is ignited in a resonator with power and size affected greatly by the design of the reactors. The primary characteristics of various different MPCVD reactors are the resonant geometry and the coupling of the microwave into the reactors. With proper design of the reactor in an MPCVD, the growth rate of

2.4.2 Growth mechanisms of CVD diamond

FIG. 2.4.1

C 0 1.0

167

Schematic of the Bachmann triangle dia-

gram [3].

M: Methanol E: Ethanol

Amorphous carbon

0.5 CO

on

arb nd c

mo

-dia

Non

CO 2

d

n mo

E

C CH4 M:E=7:3 D B M

Diamond

) [0] ])+ /([C

[H ]/([ H] )+[ C] )

] [C

0.5

dia

No

th

grow

A 1.0 H0

0 H 2O

0.5

1.0 O

[0]/([0])+[H])

diamond could be enhanced while the material uniformity and quality are improved. A number of MPCVD system designs have been proposed, such as cylindrical resonant cavity reactor, a bell jar reactor, a multimode noncylindrical cavity reactor, and an ellipsoidal cavity reactor; some of these reactors have already been produced commercially. As the gas species pass through the electric discharge activation region, they are ionized and dissociated by the microwave energy into reactive radicals and atoms. In the plasma, the external energy couples directly to free electrons, producing hydrogen radicals via H2 + e ! 2H + e The H radicals then react with neutral carbon sources such as CH4 molecules to produce reactive carbon containing radicals H + CH4 ! CH3 + H2 Recombination of the methyl radicals could induce species with more carbon atoms such as ethyl radicals CH3 + CH3 + M ! C2 H6 + M where M is a third body. Methyl radicals and ethyl radicals are believed to be the most abundant and most stable carbon-containing radical species under the MPCVD conditions. Finally, the reactive radicals are transported by diffusion and convection to the growing diamond surface. (100) and (111) faces are the two prominent low-index surfaces of diamond; unreconstructed (100) diamond has two dangling bonds per atom at the surface while the (111) surface only has one. The presence of dangling bonds means that the surfaces can be easily hydrogenated or reconstructed. Valence band and core level photoemission spectroscopy (PES), Auger electron spectroscopy (AES), low-energy electron diffraction (LEED), lowenergy electron loss spectroscopy (EELS), and photon stimulated ion desorption (PSID) are used to discover the reconstruction of surface structure. The results showed that the (100)

168

2. Semiconductor diamond

surface of diamond has a 2  1: H reconstructed structure with carbon dimer rows while the (111) surface has an unreconstructed structure with hydrogen termination. The elementary step of the growth process is the incorporation of a carbon atom into the surface. In the chemical reactions relevant to diamond growth, the hydrogen atom plays a key role. At the growth surface, the hydrogen atom can abstract hydrogen from the filled C–H sites to create free surface sites as the bond energy of the H-H bond is greater than that of the C–H bond. This reaction can also dehydrogenate the adsorbed carbon species, benefitting the incorporation of carbon atoms into the crystal lattice. The surface radical sites created by the interaction between the surface and atomic hydrogen will either be refilled with H atoms or react with carbon-containing radicals, which results in adsorption of the carbon species. Besides, the H radicals terminate the dangling bonds, stabilize the sp3 coordination of the diamond lattice from rearranging to the sp2 graphitic form, and react with any sp or sp2 to form carbon, etch the graphite phase away on the surface such that only diamond continues to grow. Atomic oxygen plays a similar role as atomic hydrogen in the removal of defects in diamond at the lower temperature range [5]. Further insight into the complex CVD growth process of reacting species and surface stereochemistry is quite difficult. Lombardi et al. showed that up to 28 species and 131 reactions could exist in a simple H2-CH4 discharge [6]. Experimental diagnosis of the plasma through the utilization of laser absorption spectroscopy and optical emission spectroscopy (OES) has revealed some aspects of the reacting species. Laser absorption spectroscopy is used to determinate the absolute column densities of stable species in plasma [7]. Meanwhile, OES measures the emission from excited states of species, providing information about the plasma temperature and electron density [8]. By those measurements, the source gas is found to be converted into a mixture dominated by CH4 and C2H2, depending on the process conditions and design of the reactor. As for reactions on the substrate surface, however, due to the presence of plasma, an atomic scale in situ observation of diamond is scarcely possible. Theoretical calculations were thus proposed as a complement to understand the growth mechanism. Semiempirical molecular orbital methods [9], Monte Carlo models [10], and quantum mechanical and hybrid quantum mechanical/molecular mechanical (QM/MM) cluster models [11] have been used to investigate the CVD growth process of diamond. It should be noted that the growth mechanisms differ on growth planes with different crystallographic orientations as the vacant sites on the surface vary [12]. The {100} surface of diamond has been a longstanding topic for both theoretical and experimental studies due to atomically smooth morphologies and the low defect density of this low-index surface. A preliminary understanding has been achieved in the explanation of how carbon-containing radicals are added to the (100) diamond surface and migrate on the surface [11, 13, 14] during the CVD growth. On the other hand, growth on the (111) surface of diamond is not as well addressed as the (100) surface [15] as the growth of highquality (111) diamond is much more difficult. However, an increasing interest has emerged in diamond (111) surface growth as the mobility of electrons in phosphorus-doped diamond (111) films is higher than that of (100) films. Therefore, more efforts should be paid to the studies of the diamond (111) surface. Even though lots of theoretical speculations have been proposed to explain the CVD growth process, the mechanism of diamond growth is still not understood thoroughly because of the large number of experimental parameters and the difficulty in experimental measurements.

2.4.3 High-speed growth of CVD diamond

169

2.4.3 High-speed growth of CVD diamond High-speed growth is one of the most important issues in the fabrication of single-crystal diamond (SCD) by MPCVD for industrial production and application. Among the methods to improve the growth rate, nitrogen addition to the reaction chemistry shows a significant effect. Yan et al. demonstrated a high growth rate of up to 150 μm/h in 2002, two orders of magnitude higher than the previous data in the literature [16]. The growth rate was further improved to 165 μm/h by Liang et al. [17] Although nitrogen inclusion is beneficial to a high growth rate, as reported in these works, defects related to the nitrogen impurity are also introduced. With a high nitrogen doping level, a dark brown discoloration appears due to the high absorption of impurities associated with nitrogen. Lowering the proportion of nitrogen addition to a few ppm can still increase the diamond growth rate because of the catalytic effect of nitrogen [18]. A 10 ppm nitrogen addition could increase the growth rate by more than a factor of two while an uncolored “optical-grade” CVD diamond can be fabricated [19]. However, though nitrogen-doped diamond can be optically transparent, nitrogen incorporation still drastically deteriorates the electronic properties of the optical-grade diamonds. Thus, nitrogen addition can only be used for optical and thermal applications of diamond. For the electronic usage of diamond, nitrogen should be avoided as much as possible. To achieve a growth environment without nitrogen addition, nitrogen impurities in the feed gases and from any leaks in the vacuum system should be eliminated. With an ultrahigh vacuum-compatible chamber, high-purity grade input gases, and a purifier, the nitrogen concentration in the high-quality diamond could be controlled to less than 0.6 ppb [20]. An increase of the methane concentration in the input gases, which will lead to an increase of carbon-containing radical density, can also boost the growth rate. However, under constant chamber pressure, the H concentration remains unchanged while the CH3 radical density is increased, which will deteriorate the crystal quality. Also, if the methane concentration is raised to more than 7%, soot formation will occur, which may cause heating and damage to the reactor quartz window [21]. Therefore, the achievement of high-speed growth by increasing the methane concentration is not a mainstream method. On the other hand, with the extremely low methane concentration of 0.025%, high-quality atomically flat diamond surfaces with a mean roughness of 0.04 nm can be obtained [22]. Though the growth rate is too slow (20 nm/h) for any practical deposition, this recipe could be used as the initial growth process before diamond deposition at a normal growth rate [23]. Substituting of hydrogen by argon in the source gas can also enhance the CVD growth rates of single-crystal diamond. This enhancement can be attributed to the increase of the plasma temperature as the thermal conductivity of argon is 10 times lower than that of hydrogen. By using optical emission spectroscopy and plasma simulation, the gas temperature is calculated to be increased by at least 500 K when argon is utilized in the plasma [24] in the CVD growth of diamond. The addition of argon to the CVD source gas will lead to a more confined and hotter plasma core, which promotes the dissociation of H2 and carbon-containing species. Therefore, the argon addition is efficient to improve the growth rate without additional adjustment of the reactor. Bolshakov et al. increased the growth rate by a factor of 2–4 with a 20% Ar addition, up to 105 μm/h at 15% CH4 [25].

170

8% CH4

0 10

10

4

2. Semiconductor diamond

/h µm

X sp2 µ G/[H]2

1E-8

6% CH4 4% CH4 2% CH4

10

10

1

Combustion Torch

/h

µm

1E-9

DC Arc Jet

1

/h

10 –2

µm

1E-10

r ba

ba

m

r

ba

r

15

50

0 m ba r

ba

m

m

ba

r

r

1E-9

m

0

0

m

0 25

20 10

1E-10

0 30

RF Torch

Hot Filament, Microwave

1E-11 1E-11

10 –4

1 0.

[CH3] (mole/cm3)

/h µm

1% CH4

2

1E-8

1E-7

1E-6

3

[H] (mole/cm ) FIG. 2.4.2

Process map in [CH3]-[H] space showing the operating ranges of the main CVD diamond growth processes [28].

The gas pressure can influence both the plasma gas temperature and the plasma volume. The theoretical calculations predict a continuous increase in growth rate with pressure [26], with the growth rate of 1 mm/h being supposedly achieved above atmospheric pressure. Silva et al. have developed the numerical models to compute the concentrations of H and CH3 in the gas phase as a function of deposition parameters, and the operating ranges of the main CVD diamond growth processes are plotted on the process map from Goodwin [27] in CH3-H space, as shown in Fig. 2.4.2. The increase in pressure can also result in a higher concentration of CH3 radicals. Meanwhile, the concentration of atomic hydrogen increases with pressure much faster than CH3 radicals. As the defect density is inversely proportional to atomic hydrogen density, by increasing the pressure in the growth chamber, the growth rate and crystal quality of the synthesized SCD can be increased at the same time [28]. Muehle et al. extended the growth pressure to 400 Torr by retuning the reactor geometry [29]. The absorbed power density is then increased to 670 W/cm3, leading to the growth of SCD at 51 μm/h with nitrogen incorporation levels below 160 ppb. When the nitrogen incorporation level is controlled to between 400–500 ppb, SCD was synthesized at a growth rate of 75 μm/h under 300 Torr with the synthesized quality of type IIa or better [30]. The high-pressure synthesis method is able to synthesize high-purity and high-quality SCDs at high rates over large deposition areas [31]. Therefore, it is currently the most prominent CVD synthesis method for high-speed growth.

2.4.4 Large-area growth of CVD diamond

171

2.4.4 Large-area growth of CVD diamond In spite of its extraordinary properties, diamond’s potential applications are restricted by the missing availability of large-area crystals. For homoepitaxial growth, the final area is limited by the diamond substrates while the most common HPHT diamond crystals are rarely above a few tens of mm2 in area. To enlarge the surface, several methods were developed, as described below: (1) High-rate repetitive growth: As the growth occurs both on the horizontal faces and vertical faces of the initial diamond seed, enlargement of the surface can be achieved by vertical growth. By combining a specific substrate holder design with nitrogen addition during growth, Makuno et al. repeated the deposition run on the same substrate many times to obtain a large CVD SCD. Large diamonds of 3.5–4.65 ct had been grown on a 27–37 mm2 seed after 24–31 times growth [32]. However, there are some major drawbacks with this repetitive growth process. The main restriction is a thick polycrystalline diamond (PCD) rim grown around the SCD substrate, limiting the further enlargement of the SCD area. In addition, the intentional nitrogen addition to improve the growth rate in this process reduces the quality of the diamond. (2) Side-surface (100) growth: Due to the difficulty in high-rate repetitive growth, a sidesurface growth method has been developed. First, a thick diamond epilayer is grown by repetition growth on the initial substrate. Second, the side surfaces normal to the previously grown one are cut and polished, then used as substrates for the next growth. Finally, the process repeated and ultimately enlarged the diamond surface. This produced a half-inch (12.6  13.3  3.7 mm) SCD [33]. However, this method is time consuming, which makes it unsuitable for practical use. More importantly, a high number of dislocations and defects exist along the interfaces between different growth surfaces, deteriorating the overall quality. (3) Mosaic wafer growth: Mosaic wafer growth is a popular approach to enlarge the CVD diamond area. It was first proposed by Geis et al. in the 1990s, with a continuous 20-μm SCD film grown on top of 20–30 oriented crystals [34]. However, early mosaic wafers usually had pin holes at the corners of each connected substrate and obvious boundaries between the neighboring substrates. Yamada et al. introduced an improved method to fabricate mosaic wafers with a smooth surface and almost invisible boundaries, as shown in Fig. 2.4.5 [36]. The key point of this improvement is the use of a so-called lift-off process. In the mosaic growth process, the CVD diamond epilayer is separated from the HPHT substrate by the lift-off process, which was developed to minimize the losses in separating [37]. To be specific, the substrate was ion-implanted first before CVD growth to form a well-defined damaged layer below the surface at the depth of a few micrometers. During the growth, high temperature causes the transformation of the implanted layer into the graphitic phase. The graphitized material can be easily removed afterward by electrochemistry etch [38]. With the use of the lift-off process, thin SCD “clones” with similar crystallographic characteristics as the “mother” substrate were fabricated. Then the clone process is repeated to form the required number of clones, as shown in steps 1–4 of Fig 2.4.3 [39]. Each “clone” plate had the same “mother” substrate. Afterward, the clones were aligned with each other to form a large substrate (step 5 in Fig 2.4.3). A diamond epilayer was then grown on the tiled clone wafer. By utilization of the lift-off process again, a freestanding large-area SCD wafer can be produced (step 6 in Fig. 2.4.3).

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2. Semiconductor diamond

FIG. 2.4.3

Schematic of the mosaic wafer growth [39].

The size of the mosaic wafer was extended to 40  60 mm2 in 2014, comprising 24 SCD clones [40]. A critical factor in the mosaic wafer growth method is the precise alignment of clones during the formation of the tiled wafer. Any misalignment could result in dislocations and stress at the interfaces. Due to the high defect density, the boundary of the joined region should act as the “dicing lines” to avoid devices there. The unusable area may be very high if the size of the devices does not fit well with the SCD plate [41]. Besides, the mosaic wafer growth is a multistep technique involving a complicated cloning procedure.

2.4.5 Rimless growth with expanding surface While progress has been made in the methods mentioned above, problems such as a long growth time, limited growth size, crystal stress, and dislocations still restrict their practical applications. A rapid and efficient growing method is require for area expansion in SCD. The rimless growth of diamond with an expanding surface has therefore come into notice. During a conventional CVD growth process, the SCD area of diamond shrinks due to the growth of a PCD rim around the substrate. Except for a reduction of the SCD area, the PCD rim also induces defects on the surface, especially at the edges. Therefore, the PCD rim should be minimized during growth while the lateral growth is promoted to expand the surface area. A growth recipe was therefore developed by Nad et al. [42] that controls the substrate temperature in a function of growth time. This recipe alternately changes the temperature, benefitting the fast growth of different crystal directions, leading to increased lateral growth to smoothen and enlarge the top surface. The SCD substrates grown via this recipe have an enlarged area of approximately two times during one continuous run. The area of the rimless diamond surface could be further increased by utilizing multigrowth steps [35]. Without the PCD, stress in the synthetized SCD can also be reduced considerably, improving the overall quality. To achieve the rimless growth, a pocket holder design is also essential apart from the growth recipe. The effect of holder design will be discussed later in this section.

2.4.6 Design of substrate holder for the growth of CVD diamond Yamada et al. theoretically investigated the influence of substrate geometry and discharge position on the synthesis environment and results [43]. They found that the

2.4.6 Design of substrate holder for the growth of CVD diamond

173

FIG. 2.4.4 The open and enclosed substrate holder designs where “d” indicates the distance between the seed and the top of the holder [44].

W

d

FIG. 2.4.5

Pocket holder design for SCD synthesis [35].

distributions of electric fields, power density, temperature, and gas flows are strongly modified by the design of the holder. Besides, qualitative correspondences between the numerical simulations and experimentally observed surface morphologies are found. In other words, the diamond synthesis process can be modified by varying the local substrate holder design. A molybdenum substrate can be broadly classified as “open” and “enclosed” types, as shown in Fig. 2.4.4. Within the corresponding factors, substrate temperature distribution brings the most important influence on the synthesis process. It was observed from the simulations that temperature distribution over the substrate is more uniform in the enclosed holders. Therefore, diamond epilayers grown in the enclosed holders have a smoother surface than those grown in open holders [45]. For enclosed holder designs, the absorbed power density concentrates on the edges of the holder, resulting in the separation of the high power density area with the substrate surface [46]. Hence, the depth of the substrate surface to the holder surface beings great influence to the crystal quality and growth rate. With a high holder depth, the beneficial effect of the high power density could not be utilized. Nad et al. developed the “pocket holder” as a kind of enclosed holder, as shown in Fig. 2.4.5 [35]. This configuration creates a thermally uniform environment that can shield the intense microwave plasma and any microwave field concentration on the edges of the substrate. The diffusion and convective flows of the plasma and the radical species onto the substrate surface can be adjusted and controlled by varying the pocket dimensions, specifically “d” and “w,” depending on the growth time. Under the pocket holder, the SCD can grow both vertically and horizontally, forming a smooth top surface without any PCD rim [30]. Through a birefringence imaging measure, the rimless surface grown in a pocket holder showed to be a stress-free, high-quality SCD [28].

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2. Semiconductor diamond

2.4.7 Low defect growth of CVD diamond Extended defects such as dislocations and stacking faults are common in both natural and synthetic diamonds. In the MPCVD diamond, a typical dislocation density is between 104 and 106 cm2, several orders of magnitude higher than the best HPHT diamonds [47]. Dislocations can generate energy levels in the bandgap of diamond due to the presence of carbon dangling bonds in the core [48]. Besides, the stress field around dislocations is high enough to cause birefringence, plaguing the use of diamond for optical windows or Raman lasers [49]. For power devices, dislocations lead to an increase in leakage current and to a reduction in the maximum breakdown field [50, 51]. Therefore, it is urgent to develop strategies aimed at reducing dislocations. Many studies have been carried out to confirm the type and propagation of dislocations in diamond by birefringence microscopy, cathodeluminescence (CL), or X-ray topography (XRT) [52, 53]. The extended defects in CVD crystals have two main origins: (i) new dislocations generated at the interface between the substrate and the epitaxial layer [54], and (ii) dislocations extended from the substrate [55]. Surface damages and containments on the substrate would induce the generation of new defects. Kato et al. showed that no new dislocations would be formed with an ultraflat substrate [56]. Numerous pretreatment methods have been developed to obtain high-quality, damage-free surfaces, as conventional mechanical polishing is difficult to control precisely. Kubota et al. induced a two-step polishing technique [57], consisting of a rough processing of 2 h with an iron plate and a fine polishing of 3 h with the addition of a hydrogen peroxide solution. The surface roughness after the process can be improved to an atomic scale smoothness. Besides, a novel polishing process utilizing ultraviolet irradiation was developed by the same university [58]. This process combines mechanochemical polishing with an ultravioletinduced photochemical reaction. The surface roughness could reach 0.2 nm Ra within 1–3 h at a high removal rate. On the other hand, oxygen plasma or hydrogen plasma can be used to etch the surface-damaged layer of diamond. After plasma etching, a macroscopically flat surface can be formed with an almost complete disappearance of polishing-induced features [59]. However, a number of etch pits appeared on the surface, indicating intrinsic defects existing in the HPHT substrate. Therefore, a further polishing combined with ion etching–inductively coupled plasma etching or chemomechanical polishing was performed to remove etch pits [60]. With a proper surface pretreatment procedure, high-quality CVD SCD with a similar defect density as the initial HPHT substrate can be obtained. To avoid extended defects from substrates, a reduction of dislocation density in the HPHT substrates is required. High-quality HPHT diamond with a dislocation density of less than 50 cm2 was fabricated by Masuya et al. [61] Mokuno et al. demonstrated SCD plates with a dislocation density below 400 cm2 by a lift-off process on the HPHT substrate with a dislocation density below 100 cm2 [62]. However, the low-dislocation HPHT diamond required a specific growth technique with slow growth rates and high cost. Growth strategies based on a standard Ib HPHT diamond are required due to the poor availability of high-quality HPHT diamonds. By using off-axis substrates with an angle higher than 10 degrees, the propagation direction of dislocations can also be changed because the <110 > direction becomes more energetically favorable [63]. Tallaire et al. achieved a reduction of dislocations in SCD by lateral growth over a macroscopic hole drilled with a laser on a standard HPHT Ib substrate [64]. Dislocations were found to propagate with a [001] or [010] direction so that they could

2.4.8 Characterization and physics of diamond

175

terminate either on the substrate surface or at the side facets before merging. The laterally grown region thus contained a low dislocation density of around 2  103 cm2. Another attempt is to engineer the substrate to a pyramid shape, which would disappear after a certain thickness [65]. The inclined faces of the pyramid deviated dislocations toward the edges of the crystal to limit their occurrence at the surface. The density of etch pits in the center was found to be around 3 103 cm2 while it was at least an order of magnitude lower around the central region [66]. Nowadays, the development of diamond-based electronic devices is strongly hampered by the lack of low-dislocation single-crystal material. Although the reduction of defects is a great challenge, recent works have constituted a first step toward achieving dislocation-free diamond crystals.

2.4.8 Characterization and physics of diamond The research on the carrier dynamics in diamond, the scattering mechanism of carriers, and the physical properties of excitons and other basic materials have been the research focus of scientists. The research results have profoundly influenced the quality of material growth and the device development process. In diamond, room-temperature electron and hole mobilities of 2000 cm2 V1 s1 have been reported in most of the literature. In 2002, Isberg from the University of Uppsala in Sweden measured room-temperature drift mobilities of 4500 cm2 V1 s1 for electrons and 3800 cm2 V1 s1 for holes in high-quality single-crystal diamonds using the time of flight method [1]. Moreover, advanced time-resolved cyclotron resonance measurements have shown low-temperature mobilities exceeding 106 cm2 V1 s1 and the evaluated mobility exceeding 5000 cm2 V1 s1 for holes and 14,000 cm2 V1 s1 for electrons. The high mobilities that have been achieved are of great advantage to other semiconductor materials, making them potentially useful in electronic devices in the near future. The temperature would have influence on the mobilities, which is a research spot. Several research groups have contributed to this work. Isberg systematically studied the carrier transport mechanism in the temperature range of 80–470 K, and obtained the low field hole mobility (2600–3800 cm2 V1 s1). By changing the intensity of optical excitation over several orders of magnitude, it is shown that the space charge effect has an impact on the mobility. It is pointed out that acoustic phonon scattering is the dominant scattering mechanism at temperatures below 350 K [67]. In 2009, Gkoumas from the University of Surrey found that the transient current pulses of hole mobility increase strongly at reduced temperatures, which is consistent with acoustic phonon scattering processes. However, electron mobility values appear to remain relatively constant with lower temperatures, suggesting that different mechanisms than optical or acoustic phonon scattering are limiting the charge transport [68]. In 2011, Isberg systematically studied the carrier mobility at a temperature of 83–460 K and for electric fields between 90 and 4 103 V/cm. It was found that in the low temperature and low field, hot electron transfer effect occurred. The influence of temperature on electron mobility shows that acoustic phonon scattering plays a dominant role at lower temperatures while at higher temperatures, interstellar phonon scattering dominates [69]. Besides temperature, the electrical fields would affect the mobilities, too. Pernegger from Switzerland further studied the effect of electric field intensity (0.2–1.5 V/μm) on carrier drift

176

2. Semiconductor diamond

velocity and mobility [70]. In 2006, Isberg used the transient photocurrent method to study the distribution of the electric field at different depths of the crystal, and proposed that the space charge accumulation has an impact on the electric field [71]. Later, they observed the phenomenon of negative differential mobility that is related to the electron regeneration between different energy troughs in the conduction band below 140 K by measuring the drift velocity of electrons in diamond as a function of the applied electric field [72]. After different processes, the diamond surfaces could show oxygen termination or hydrogen termination. Carrier transport would behave different under different conditions. Researchers at Hasselt University studied the transport of diamond carriers and the properties of light emission for completely oxidized and completely hydrogenated. They found that the oxygen termination will generate more surface states, which will affect the carrier transport and luminescence properties [73]. In addition, they have also studied the impact of different electrodes (such as Al, Au, etc.) on the carrier transport properties of hydrogenterminated or oxygen-terminated diamond [73, 74]. For ultrathin diamond films, the electrical transport would be different from the normal state. Isberg first set up a lateral ToF charge carrier transport measurement system. The measured near-surface hole drift mobility is reported to be about 860 cm2/V
s across a contact spacing of 0.3 mm, which shows the influence of the surface state on the carrier transport [75]. Because diamonds have some extreme properties, some interesting phenomena could also be observed in this material. The transient-current technique can also be used to detect small concentrations of charged defects in a diamond and to study its photoionization spectrum. By continuously measuring the charge concentration while illuminating the samples with monochromatic light, the evolution of the charge state of the dominating defect can be continuously monitored. A photoionization cross-section spectrum from the dominant deep defect, which can be attributed to the single substitutional nitrogen impurity, can be obtained [76]. The generation, transport (across macroscopic distances), and detection of valley-polarized electrons in bulk diamond with a relaxation time of 300 ns at 77 K was also discovered in 2013 [77]. Despite its outstanding electronic properties, diamond also holds excellent optical properties. Impurity-free diamond has high transmittance from the UV to the far IR band, which makes it an ideal window material for high-power lasers and detectors as well as protective films for optical lenses. Most of the properties of diamond will change when there are impurities, especially the optical properties. For example, diamond without impurities is colorless while diamond with nitrogen would exhibit yellow, green, black, etc. The appearance of these colors depends on the type, presence, and concentrations of the impurity elements in the diamond. The research on the optical properties of diamond has been conducted since last century. In general, CL and photoluminescence (PL) are the two major techniques for the optical characteristics of diamond. Most research on band-edge emission in diamond has been studied by CL while less has been studied by PL because optical excitation is difficult due to the bandgap width of diamond. However, an electron beam will cause a certain degree of damage while PL measurements are nondestructive to diamond samples. In addition, charge accumulation would happen on the diamond surface when CL is performing.

2.4.9 Prospective

177

The intrinsic and extrinsic recombination radiation from natural and synthetic aluminumdoped diamond were reported by Dean in 1965. The bound excitons were reported to have a thermal and optical ionization energy of  50 meV [78]. Free-exciton recombination have been observed in polycrystalline diamond or in nitrogen-induced samples by Kawarada from Waseda University [79]. The temperature has much influence on the emissions of both free and bound excitons. Generally, the intentions of the PL/CL spectra usually decay with increasing temperature. Researchers from Japan reported that free-exciton emission of high-purity chemical vapor deposition (CVD) diamond keeps almost constant at a certain temperature range (80–170K), but was reduced by two orders of magnitude at room temperature [80]. Through temperature-dependent measurement of optical properties, more physical mechanisms behind can be obtained. Temperature dependence on free excitons can be analyzed using a rate-equation model including a bound exciton formation. Moreover, the lifetime of indirect free excitons can be estimated [81]. Researchers from Kumamoto University in Japan first reported the quantum efficiency η of the luminescence and the radiative transition probability of the indirect excitons by measuring the time decay of the luminescence and the absolute intensities of incident, reflected, and transmitted light of the excitation laser at 83 K. The estimated values are 2.27 102 and 4.3 105 s1, respectively. And the radiative lifetime was reported to be 2.3 μs [82]. The longest decay time of 80 ns was measured around 100 K in nondoped high-temperature high-pressure single-crystal diamond samples through transient measurement of excitonic luminescence under pulsed excitation [83]. The existence of shallow traps can contribute a lot to the optical characteristics. For example, it could decrease the free-exction intensity at a certain temperature range [84]. However, everything has two sides. The existence of some specific trap-vacancy centers, such as nitrogen-vacancy and silicon-vacancy centers, has a huge potential application in quantum computing and quantum information processing due to its good controllability, high quantum efficiency, uniform line width, and long decoherence time. The nitrogen-vacancy center (NV center) is one of numerous point defects formed by coupling one substitutional nitrogen atom and an adjacent vacancy. The photoluminescence characteristics of NV centers have been deeply explored, especially in the negative charge state (NV). Electron spins can be manipulated by applying a magnetic field, electric field, microwave radiation, or light at room temperature. The unique optical characteristics were caused by the energy level structures, which can be described by group theory. The NV center system in diamond is considered to be one of the promising candidates for quantum information processing and quantum computing in the future, and remains a hot research spot these days.

2.4.9 Prospective Thousands of years ago, the materials of diamond were discovered by humans. Inspired by its sparkling appearance under the sunlight, diamond has long been known as a part of highend jewelry. With the gradual discovery of diamond’s various properties such as mechanics, electronics, and optics, people have developed a keen interest in its use in other fields in addition to jewelry. Due to the scarcity of natural diamonds and their high price, most natural diamonds are used as raw materials for jewelry processing. Since the discovery that diamond

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2. Semiconductor diamond

is the hardest substance in the world, people began to consider diamond as a cutting tool. With the invention of high-temperature and high-pressure technology in the 1950s, synthetic diamond has been widely used in industrial production such as drills and grinding tools. However, the application of diamond in the semiconductor field began in recent decades. In the early days, the progress of the properties of semiconductor-grade diamond materials was not very satisfactory. When the chemical vapor deposition method was adopted, diamond research revived. The research direction of diamond crystal growth gradually shifted from polycrystalline to single crystal. Constrained by the lack of early research direction and experimental equipment, the semiconducting properties of diamond materials were far from the desired results. To realize the theoretical expectation of the performance of diamond materials, the growth of single-crystal diamond and the control of semiconductor electrical properties are the key elements. At present, the high-quality single-crystal diamond is obtained by microwave plasma chemical vapor deposition. In this method, the substrate is made of millimeter-scale high-temperature and high-pressure diamond material. In order to realize the preparation of a large diamond, it is expected that the high-temperature and high-pressure method will continue to develop within a period of time in the future, and it will eventually be able to provide inch-level substrates. At the same time, microwave plasma chemical vapor deposition will still be the preferred method for the preparation of high-quality diamonds in the near future. There will be major breakthroughs in the simulation of growth processes, innovation of reactants, improvement of growth processes, and transformation of deposition equipment. At present, the highest growth rate of 165 μm/h in the world cannot meet the application requirements of semiconductor-grade diamond materials. It is expected that there will be a substantial increase in growth rate in the future, reaching millimeters and even centimeters per hour. In terms of crystal size, the size of diamonds prepared at present is in the order of centimeters, and the method of increasing the size mainly adopts a repetitive growth method, a three-dimensional growth method, or a splicing method. These methods could increase the size of the diamond to some extent, but it is difficult to substantially increase the size. It is expected that there will be a revolutionary breakthrough in diamond growth methods in the future, finally realizing the preparation of inch-scale diamond substrate materials. For different application requirements, the crystal quality improvement of diamond has different directions. Therefore, in the future, the improvement of the crystal quality of diamond will make appropriate improvement according to different improvement goals. However, it will generally develop in the direction of high quality, high reliability, and large scale. And, a new perfect diamond crystal quality standard certification system will appear.

Acknowledgments This project was supported by the National Key Research and Development Program of China (Grant No. 2018YFB0406501), the Beijing Municipal Science and Technology Commission (Grant No. Z181100004418009).

References [1] J. Isberg, J. Hammersberg, E. Johansson, T. Wikstrom, D.J. Twitchen, A.J. Whitehead, S.E. Coe, G.A. Scarsbrook, Science 297 (5587) (2002) 1670–1672.

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2.5.1 Introduction

181

C H A P T E R

2.5 Single-crystal diamond wafer production Guangchao Chen College of Materials Science and Opto-Electronic Technology, University of Chinese Academy of Sciences, Beijing, P. R. China

With the development of CVD diamond techniques, large-bulk and/or large-area diamonds were able to grow out. In light of applications in the industrial area, wafer processing should be considered so that the diamond wafer could be obtained to support the requirement of various applications in society. Conventional wafer processing includes cropping, trimming, slicing, edge rounding, grinding, lapping, polishing, etc. Each procedure was reviewed to analyze the possibility of utilization in diamond wafer processing. The emphasized point was the polishing procedure in this chapter. By these analyses, conclusions were deduced. One of the important points was that the diamond wafer was preferred to be grown out rather than produced by conventional wafer processing.

2.5.1 Introduction Diamond is a unique material with a wide bandgap. After more than 30 years of efforts to grow diamond by chemical vapor deposition (CVD) [1] especially, the achievement of largebulk single-crystal diamond growth in 2002 [2], single-crystal diamond wafer has been expected for electronic devices fabrication [3]. Usually, the wafer is the cut slice of artificially grown crystals with a certain surface condition so that they can be used for electronic, optical, or other applications. In light of applications in the microelectronic field, a wafer is expected to be of large size with a flat and high-quality surface because the large size can reduce the cost of microelectronic device fabrication, and the flat and high-quality surface can ensure the microelectronic device performance [4]. Wafer processing mainly includes cropping, trimming, slicing, edge rounding, grinding, lapping, polishing, etc. [5] After wafer processing, the finished surface condition should be characterized by a degree of crystalline perfection; defect, microstructure, and surface/subsurface damage; global surface planarization; uniform thickness/total thickness variation (TTV); residual stresses; and roughness. The economic factor, efficiency, cost of supply material, and loss of row crystal are the main issues in wafer processing.

182

2. Semiconductor diamond

By wafer processing, an Si wafer can be 8 in. in diameter with a very fine surface [6]. An SiC wafer, another wide-band semiconductor as diamond, can be 8 in. in diameter [7]. A series of inch-graded commercial products of both the Si wafer and the SiC wafer can be bought in the mass market. By contrast, a single-crystal diamond wafer is much smaller in size, and there has not been a commercial product with a size larger than 4 in. in diameter so far. Actually, the wafer processing is difficult to apply in a single-crystal diamond, let alone a polycrystalline diamond film. For example, it is difficult to cut diamond by common wire saw, and no cheap abrasive supply material has been found for diamond grinding and polishing [8]. The problems also exist in the polishing procedure. Besides the high hardness and stable chemical property of diamond, the anisotropy is the major factor making the polishing complex and difficult. This is because it leads to nonuniformed material removal, depending on the crystal orientations and polishing directions. In addition, the existence of anisotropy requires the extra skills and experience of the worker, which reduces the polishing efficiency. Though anisotropic polishing mechanisms have been studied [9], the effect of the anisotropy on the surface quality of polished surfaces has not yet been fully understood. Generally speaking, the present polishing procedure is not easy to transfer to manufacturing the single-crystal diamond wafer. In order to find much more efficient approaches to the single-crystal diamond wafer, in this chapter, we review the techniques according to the wafer routine. As CVD technology is a possible method to obtain the large wafer, the review is focused on the CVD single-crystal diamond wafer. The contents are mainly referenced by H.-J. M€ oller, Y. Chen, H. Li, Y.N. Palyanov, and O. Ohnishi [ 4, 8, 10–12].

2.5.2 CVD growth of single-crystal diamond wafer Single-crystal diamond can be grown homoepitaxily or heteroepitaxily. It has been reported that both homoepitaxical growth and heteroepitaxical growth can adopt the microwave plasma-assisted chemical vapor deposition (MW CVD) technique and the hot filament CVD technique [2, 13]. DC Arcjet CVD has been used in homoepitaxical growth rather than in heteroepitaxical growth [14]. Both MW CVD and DC Arcjet CVD can grow millimeter-gradesized diamond, or even larger. It has been reported that a 4-in. single-crystal diamond wafer was grown by the MW CVD technique [15], and a 7 7 mm single-crystal diamond wafer was grown by DC Arcjet CVD [16]. The following content would specify each technique in singlecrystal diamond growth.

2.5.2.1 MW CVD technique [10, 11] MW CVD possesses the highest growth rate of more than 150 μm/h [2], and a large area over 90 mm (nearly 4 in.) in diameter at present [15]. However, it was once called the “low growth technique.” Since 1983, when the first report about the MW CVD of diamond films was presented, its growth rate has never been higher than 10 micrometer/h, and the deposited area no more than 1 cm2 until the end of the last century. However, after 20 years of effort, things changed in 2002 when a large bulk single-crystal diamond was successfully grown at a growth rate of more than tens of micrometers/h [2]. Then, the repetition high-rate growth

2.5.2 CVD growth of single-crystal diamond wafer

183

method, the mosaic growth method, and the “clone” with lift-off process were gradually proposed to produce wafers. In these improved MW CVDs, the pressure in the chamber (or reactor) was regarded as an important factor to influence the growth rate. It was speculated that the growth rate for a single-crystalline CVD diamond would increase with the chamber pressure, and might arrive at 1 mm/h under pressure above 1 atm based on the theoretical calculation. One possible reason was that the temperature of plasma increased with the increase of the pressure in the chamber (or reactor). For example, the gas temperature in the plasma was about 2800 K @100 Torr and 5 kW input MW power with 2.45 GHz, and it increased with the pressure increasing. The high temperature of plasma would result in not only the high density of radicals for deposition, but the fast surface reaction on the substrate [17]. Another factor to influence the growth rate was the addition of N2 in the deposition. It was found that the addition of N2 in the deposition area could increase the growth rate dramatically. For example, a few ppm of N2 concentrations in the gas phase could strongly increase the crystal growth rates on <100> orientations due to a catalytic effect. Although the synthesis of millimeter-thick crystals was often performed with the addition of several hundreds of ppm of N2, it should be used carefully for the addition of N2 because nitrogen in diamond deteriorates the electronic properties. For electronic application, N content should be at the level of a few ppb in the crystal. To increase the deposition area, the plasma-ignited frequency has been selected because the plasma ball diameter (D) was approximately half the radiation wavelength, that is, D 12 λ. It would be 6.1 cm @ 2.45 GHz and 16.4 cm @ 915 MHz. So, 915 MHz was adopted instead of the usual 2.45 GHz. On the contrary, an absolute inverse thought was proposed to increase the deposition by millimeter waves, the shorter waves than the microwave. For example, the flat plasma excited by a 30-GHz and 28-GHz gyrotron source was used to deposit film of 60– 90 mm in diameter or film extending over 100 mm, respectively. In these cases, polycrystalline diamond films instead of single-crystal diamond wafer were deposited with a growth rate of about 10–15 μm/h. Fig. 2.5.1 is the MW CVD instrument for polycrystalline growth in the past and the single crystal at present, respectively [17, 18]. The growth parameters are listed in Table 2.5.1, corresponding to each growth procedure [19, 20]. It can be seen that the chamber becomes short and wide. The distribution of electromagnetic energy is shown in Fig. 2.5.2 [18]. Therefore, the electromagnetic energy is concentrated in a wide area with short height. As the flux of work gas is several sccm order of magnitude, the fluid effect on the growth can be ignored. This means that the shape of the electromagnetic field is the key factor in MW CVD. The morphology of the as-grown single-crystal diamond always possesses a relatively smooth top terminal face with a quite rough lateral surface due to polycrystal growth. This phenomenon is serious in high growth rate repetition growth (including 3D growth). As the thickness is not very large, the lateral surface is not paid much more attention in mosaic growth. The morphology of each as-grown crystal is illustrated in Fig. 2.5.3 [21].

2.5.2.2 DC arcjet technique DC Arcjet CVD was first reported to deposit diamond in 1988 [22]. Fig. 2.5.4 illustrates the instrument of DC Arcjet. It was in 1990 that the highest growth rate of more than 900 μm/h (this value still remains unsurpassed by any other CVD technique) and about an 8%

Microwave Gas Inlet

Water cooling Plasma Cu cylinder

z

Substrate

z

Plasma

Heating

Quartz window

Pump

Microwave

Bias power supply 0 ± 300 V DC

Substrate

(a)

Microwave

(b)

(B)

(A)

FIG. 2.5.1 instrument of microwave CVD. (A) For polycrystalline diamond film [17]. (B) For single-crystalline diamond film [18]. TABLE 2.5.1

Comparison of growth parameters for different diamonds in MW CVD

Diamond

Power (w)

Frequency (GHz)

Gas flux (sccm)

Chamber pressure (×1000 Pa)

FCH4/FH2 (%)

Substrate temperature (°C)

Polycrystala

2600

2.45

200

11

3

700–1000

Single crystalb

1200

2.45

500

8 27

1–6

10002000

a b

Ref. [19]. Ref. [20].

(A)

Pyrometer

5×104V/m

0

Glass viewing port

(B)

Source gas

5×1018/m3

Domain for calculation

Thermocouple A B

0

Conductor Substrate Plasma Coolant

10mm MW

(C) 3×108W/m3

MW ∅ 26mm

0

FIG. 2.5.2 The improved instrument of MW CVD and numerically obtained contours of (A) the strength of the electric field, (B) number density of electrons, and (C) power density [18].

2.5.2 CVD growth of single-crystal diamond wafer

FIG. 2.5.3

185

Images of grown diamond by repetition growth. (A) Top view, (B) side view [21]. Temperature monitoring window Mobile OES fiber holder Substrate

Arcjet nozzle Plasma Cover plate Cooling stage

Substrate holder

Mobile sample stage

Chamber

(A) FIG. 2.5.4

cooling water

pump

(B)

Instrument of DC Arcjet for single-crystal growth. (A) Schematic draw [23]. (B) Real instrument in

working [16].

conversion of carbon from methane to diamond was accomplished, but the deposition area was as small as several square millimeters [24]. Attracted by the high growth rate, scientists made improvements in technique in the following years. With a very high purity of sp3 content, the wafer was fabricated to arrive at 50 mm in diameter in 1992 and 110 mm in diameter in 1998 [25, 26]. So, in the 1990s, DC Arcjet CVD was regarded as technique high growth rate technique, filling the hope for realization of an industrial product. However, MWCVD developed more rapidly than DC Arcjet CVD to grow single-crystal diamond. The first report about single-crystal diamond grain was in 2010 regarding the technique called conventional wafer processing [14]. In the following year, the authors struggled for an increase of the growth area and growth rate. By these efforts, a 7  7 mm2 growth area and more than 50 μm/h were achieved [16, 23]. However, this has not yet been able to overtake MW CVD. No matter how low the growth rate or how small the growth area is, DC Arcjet CVD offers a promising procedure to continuously grow single crystal diamond, named “reposition growth” [23]. In this procedure, the substrate is continuously descending during growth so that the growth surface is maintained in a relatively stable gas environment. The key factor in this procedure is controlling the substrate descent velocity. The proper descent velocity can

186 TABLE 2.5.2

2. Semiconductor diamond

Comparison of growth parameters for different diamonds in DC Arcjet CVD

Diamond

Power (kW)

Gas flux (slm)

Chamber pressure (×1000 Pa)

FCH4/FH2 (%)

Substrate temperature (°C)

Polycrystala

100

10

330

10

7001200

Single crystalb

20

1020

46

0.52.5

9001030

a b

Ref. [28]. Ref. [16].

FIG. 2.5.5

Morphology of as-grown single-crystal diamond by DC Arcjet [27].

enhance the single-crystal growth rate, but too high a descending velocity may result in an AT-G-type morphology [23]. Besides the “reposition growth,” the fixed substrate growth is also adopted. In this procedure, the growth parameters are almost the same as that of polycrystal, and it meets the same trouble as that of MW CVD, that is, polycrystal occurs and removal procedures have to be added to perform the following growth [27]. It is useful to point out that a multisubstrate can be adopted in the fixed substrate growth procedure so that a large quantity of product could be obtained [16, 23]. The usual growth parameters are listed in Table 2.5.2 [16, 28]. The morphology of the as-grown top terminal face is always accumulated by a circle of polycrystalline grains. As the seed was wedded to the substrate holder, the lateral surface was not very coarse but had a lot of contaminants. The morphology is shown in Fig. 2.5.5 [27].

2.5.2.3 Comparison with the present single-crystal growth techniques So far, MW CVD has been the mainstream technique to grow single-crystal diamond wafers. Repetition of the high rate growth method, the mosaic crystal growth method, and the clone with lift-off method has been developed based on the MW CVD technique [29]. DC arcjet shows another approach to single-crystal diamond growth. The superiority of this technique lies in its high enthalpy of plasma, which enables the plasma to possess much more radical density benefitting the diamond growth. It may result in a higher growth rate than that in the MW CVD technique. However, neither the grow rate nor the growth area exceeded those of MW CVD at present.

2.5.3 Single-crystal diamond wafer process [4, 8]

187

FIG. 2.5.6 The morphology of as-grown single-crystal diamond and the diffraction pattern [31].

It is worth analyzing the weaknesses and virtues of MW CVD and DC arcjet so that the tendency to obtain the wafer-grade single crystal could be realized. For MW CVD, the procedure was complex for both 3D growth and mosaic growth as well as the clone plus lift-off growth method. It is hard to produce high-quality wafers as well as the large quantities used in the electronic field. MW CVD strongly depends on the electromagnetic field distribution in the growth area. The problem to overcome is the large-size wafer growth for MW CVD. Although a 4-in. wafer was achieved, is it possible for wafers of 6, 12, even 18 in., as that of the Si wafer? For DC arcjet, the plasma flow feature is more important rather than the electromagnetic field. If proper background pressure in the growth chamber was designed and maintained, the flame and diameter of the jetted plasma would be elongated as needed. This indicated that DC arcjet was much more promising for a high growth rate and large area of single-crystal diamond deposition rather than MW CVD. On the other hand, controlling the flow field is easier than that of the electronic field. So, DC arcjet may be easy to operate. Although DC arcjet possess so many merits, the pollution of metal from the metal electrode is the big problem for DC arcjet CVD. RF CVD was first used to deposit diamond in 1987, one year earlier than that of DC arcjet CVD [30]. Planet-inductive RF offers a large area, and an RF jet offers high enthalpy plasma. An obvious merit is a clean plasma source without metal contaminants as in MW CVD. Recent works show that a dual-frequency RF ICP jet could grow single-crystal diamond, and it seemed to be suitable for large-area deposition [31]. Fig. 2.5.6 shows the morphology of asgrown single-crystal diamond and the electron diffraction pattern [31].

2.5.3 Single-crystal diamond wafer process [4, 8] Typical wafer manufacturing (or wafer production) refers to a process of producing singlecrystal or polycrystalline wafers from crystal ingots (or boules) of different sizes and

188

2. Semiconductor diamond

materials. It includes a series of processes, beginning with crystal growth and ending with prime wafers, that includes the major procedures such as slicing, grinding, lapping, and polishing. Having undergone these procedures, the treated wafer possesses the conditions of surface and bulk in requirement. Each procedure is important because the procedure decides not only the quality of the wafer, but also influences the cost of wafer manufacturing. However, these typical wafer processes have to face the difficulties in diamond wafer production. It is well known that diamond, as the hardest material in nature, plays an important role in other material wafer processing. The multiwire saw technique, the mainstream technique in the modern slicing procedure, is based on a diamond tool while in the grinding, lapping, and polishing processes, diamond works as an excellent abrasive material. Unfortunately, these tools and methods for other materials are hardly utilized in diamond wafer production. Besides, some procedures have to be added to produce a diamond wafer. For example, the polycrystalline grain must be removed first just after the crystal growth in CVD. In this section, we will introduce the processes of wafer production, targeting the CVD single-crystal diamond. The HTHP diamond wafer can mimic these processes by reducing or increasing some of these procedures.

2.5.3.1 Polycrystalline grain removal After the growth of single-crystalline diamond by CVD, the surfaces of the wafer or the bulk are usually covered by a quantity of polycrystalline grains. To start the wafer production procedure, the polycrystalline grains must be removed. Due to the crystalline anisotropy and the chemical stability of diamond, it is hard to remove these grains by simple mechanical ablation and chemical corrosion. Indicated by the methods originating from the polishing of polycrystalline CVD diamond films, the chemomechanical polishing (CMP), the thermochemical polishing (TCP), and the laser beam polishing (LBP) are utilized to remove these codeposited polycrystalline grains. As the treated surface is not a fine finish, this polycrystalline grain removal procedure can not be categorized as a “polishing” procedure, although the word “polishing” is in their names. 2.5.3.1.1 Chemomechanical polishing Chemomechanical polishing (CMP) is a technique that combines mechanical polishing (MP) with chemical etching in order to enhance the material removal rate and improve surface quality. During CMP, the diamond surface is treated by a rotating polishing plate, assisted by oxidizing chemicals at elevated temperatures. Usually, the diamond sample is exerted at the predetermined load. The polishing temperature is slightly above the melting point of the oxidizing agents so that the oxidation reaction can occur between the diamond and the molten oxidizing agents. So, CMP is also known as thermal oxidation polishing. The difference between CMP and mechanical ablation with the enhancement of an oxidation reaction is the temperature, which is determined by the added chemicals during the polishing. The material removal rate in CMP is dependent on the oxidization reagent used and the corresponding polishing temperature, the applied load (pressure), and the sliding speed. The polishing is normally conducted at a temperature above the melting point of the agent for an effective oxidation reaction. The higher temperature results in a higher material removal rate as the rate of chemical reaction increases faster at higher temperatures. However, higher

2.5.3 Single-crystal diamond wafer process [4, 8]

189

temperatures will make the polishing setup more complicated. All the chemomechanical polishing experiments show that a higher sliding speed and applied load result in a higher material removal rate. The chemical agents in CMP are selected from a group of oxidizing agents such as NaNO3, KNO3, KOH, KClO3, K2Cr2O7, H2O2, HClO, HNO3, H2SO4, AgO, Cr2O3, MnO2, BaO2, PdO2, and their mixtures. The commonly used ones are NaNO3, KNO3, and KOH. Their melting temperatures are 308, 324, and 360°C, respectively. In practice, some mixtures of oxidizing agents with lower melting points have been used in the polishing process to reduce the operating temperature and increase the material removal rate, such as LiNO3 + KNO3, which has a melting point at 130°C. Sometimes, diamond powder as an abrasive can be added in these chemical agents to form a slurry for more effective polishing. For example, a slurry containing diamond powder and chemicals of KMnO4 and H2SO4 is used in polishing at a temperature of 70°C. During a polishing operation, the polishing plate needs to be maintained at a constant temperature over the melting point of oxidizing chemicals dispersed onto its surface. A typical polishing parameter is a sliding speed of approximately 40 m/s (at 2800 rpm) and a pressure of 1.4 MPa (load 13 N on a 3  3 mm2 specimen). During CMP, the entire surface of the diamond is covered with molten oxidizing reagent, but polishing occurs only on peaks that are in contact with the polishing plate. It is very important to set up an oxidizing environment in the process of chemomechanical polishing of diamond. Usually, the polishing rate for CMP is on the order of 0.5 μm/h, and the minimum Ra could be even less than 10 nm at the local area. Moreover, chemomechanical polishing causes less surface damage than those by MP and takes shorter process time. The schematic diagram of the CMP instrument is shown in Fig. 2.5.7 [32].

Rotation movement

Application of pressure

Upper polish module

Polishing pad

Carrier-film and wafer

Mixing slurry Under polish module

Orbital-type CMP

FIG. 2.5.7

Schematic diagram of the CMP instrument [32].

190

2. Semiconductor diamond

2.5.3.1.2 Thermochemical polishing Thermochemical polishing (TCP) is another method developed for polishing CVD diamond films, and can be used to remove the polycrystal cogrowing with the single-crystal diamond by CVD. This technique is based on the thermochemical reaction between a diamond surface and a metal plate with high carbon solubility at elevated temperatures. In this procedure, two types of courses occur, that is, hot metal plate polishing and thermodiffusion etching. As no high pressure is exerted on the diamond wafer, microcracks or other defects on the polished diamond surfaces are not easily generated after polishing. For hot-metal-plate polishing, the diamond surface is pressed against a rotating transition metal plate at an elevated temperature from 700 to 1000°C. In this course, the diamond dissolves into the metal at the contact surface due to the chemical reactions of the diamond surface with the hot metal plate, including graphitization and diffusion of carbon atoms from the diamond surface into the hot metal plate. In order to prevent oxidation of the diamond film and polishing the plate at high temperatures, polishing is normally carried out in evacuated, reductive, or inert gas environments. The material removal rate and polished surface quality are influenced by the effect of polishing parameters, such as polishing temperature, pressure exerted on the diamond, surrounding atmosphere, sliding speed, and crystallographic orientations. Among these parameters, the polishing temperature and the degree of contact between the polishing plate and the diamond surfaces profoundly influence the material removal rate. There exists a critical polishing temperature, for example, 700°C, where polishing does not proceed because of insufficient chemical reactivity. A vacuum may be the most atmosphere for TCP because the polishing rate is the highest one among those obtained in the atmosphere of hydrogen, argon, helium, and nitrogen. Among those gas atmospheres, a significantly high value is obtained in hydrogen. This may be due to the effect of decarburization of the iron disk by hydrogen via methane molecule formation, thus keeping the disk in the active state. Therefore, a gas mixture comprising 4% hydrogen and 96% argon has also been used as the ambient atmosphere in polishing for safety reasons and to obtain a very fine surface. The other parameters such as pressure, sliding speed, and crystallographic orientation do influence the polishing efficiency and quality of the surface, but not too significantly. Thermodiffusion etching was developed by using the principle of diffusion reactions of carbon into carbon-soluble metals/alloys. These metals include iron, Mn, lanthanum (La), cerium (Ce), and their alloys. This technique is based on the atomic dissolution of carbon into the hot metal/alloy, the transformation of diamond into graphite, and diamond oxidation. The diffusion transfer of carbon from CVD diamond films to transition metal foils and molten or partially molten metals/alloys creates relatively smooth surfaces by eliminating the roughness from the top faceted surface of the film. Utilizing this thermochemical etching polishing method, experiments show that the material removal rate increases rapidly with higher temperatures using a selected metal/alloy. Polishing is more effective in a liquid alloy compared to the solid metal/alloys. Although increasing temperature and time can significantly improve the material removal rate and surface finish of diamond films, very high temperatures and too much time will make the surface rougher due to preferred etching at the boundary. The schematic diagram of the instrument is shown in Fig. 2.5.8 [33].

2.5.3 Single-crystal diamond wafer process [4, 8]

FIG. 2.5.8

191

Schematic diagram of the setup for thermochemical polishing experiments [33].

2.5.3.1.3 Laser polishing Laser polishing (LP) performs the ablation by focusing the scanning laser beam on the diamond surface. It originates from the noncontact polishing technique for polycrystalline diamond film. The operating parameters of the laser beam include wavelength (nm), energy density fluence (J/cm2), pulse length (nanosecond), repetition rate (Hz), angle of incidence (deg.), and spot size. The properties of the polished diamond, such as spectral absorptivity at the laser beam wavelength, thermal diffusivity, and purity (constituent phases, defects, and surface cleanliness) also influence the polishing efficiency. In LP, the photon energy of the incident laser is required to exceed the bandgap of diamond (5.5 eV) so that the ablation can take place directly by the interaction of the laser light with the diamond. Therefore, the wavelength should be shorter than 2.2 μm for the incident laser. If the wavelength was longer, the incident laser easily transmits pure diamond without any ablation. The strong absorption takes place only in lattice defects and impurities. This strong absorption results in the graphitic layer, which enhances the absorption. The widely used lasers for diamond polishing are two types of pulsed ones. The first consists of excimer lasers operating in a near-ultraviolet (UV) range at wavelengths from 193 to 351 nm. For this band, the excimer laser possesses the high optical absorption coefficient in diamond that can provide high-energy deposition in a small volume for rapid and complete ablation. The second is the Nd-YAG laser operating in the visible and near-infrared (NIR) spectral regions with wavelengths from 500 to 1060 nm. The high peak power of this band laser can rapidly smooth thick diamond films with a high value of surface roughness. There exists a threshold value of laser fluencies (energy density) for diamond polishing, dependent on the laser wavelength, pulse duration, and diamond materials. It is about 7 J/cm2 for a natural Ia-type diamond crystal. If higher than this threshold value, the polishing rate depends on laser parameters (e.g., laser power, wavelength, mode of operation), material parameters (initial surface, microstructure, and thickness), and process parameters (e.g., speed, focal plane position, frequency, energy, pulse duration, assist gas type, and pressure). Deep analysis shows that the removal rate is determined by the thermal conductivity and diamond crystalline quality because these two properties influence the light absorption of diamond.

192

2. Semiconductor diamond

A well-done PL method is proposed as multiple sequential lasers polishing the polycrystalline diamond film. In this procedure, an Nd-YAG (λ ¼ 532 nm) laser is first used to yield rapid and uniform material removal, then an excimer (λ¼ 193 nm @800 J/cm2) laser follows to produce a smooth surface by recovering the surface damaged due to the former laser radiation. By this method, surface roughness can be reduced from 30 to 1 μm in approximately 50 s. Actually, we have tried a hybrid method to smooth the CVD single-crystal diamond wafer by the combination of an excimer (λ ¼ 248 nm @J/cm2) laser and mechanical polishing. The roughness of the as-grown crystal is Ra ¼ 150 μm. By this hybrid method, it becomes 40 nm after 10 min LP and 20 min MP. LP is appropriate for small areas or can be extended to large areas. Polishing can be accomplished in air and is easily automated while not producing bulk heating. Especially, the NdYAG laser is ideally suited for rapid and uniform removal of diamond films. Although no force is directly applied to the polishing surface, the laser machining of single-crystal diamond is accompanied by the formation of microcracks and cleavages. These are induced by the laser pulse itself and by the volume expansion involved in the diamond-graphite transition, resulting in a rough-edged cut surface. Therefore, LP is well suited for coarse polishing of rough polycrystalline diamond films. That is why LP is hereby introduced for polycrystalline grain removal. The instrument of LP is shown in Fig. 2.5.9 [34].

2.5.3.2 Cutting diamond In the conventional procedure, cutting is the first step to slice the ingot. As there is no large enough ingot of diamond available, the cutting procedure changes from the procedure of chopping the ingot to the procedure of trimming the wafer. It means that cutting is done to obtain the regular wafer shape. In light of this point, reducing the cutting loss becomes the main issue. Obviously, the modern multisaw technique cannot handle this job. The reason is not only that the multisaw based on the diamond abrasive is difficult to saw the diamond itself, but also the kerf, 0.5 mm, is too huge for the epitaxial diamond layer to lose the grown material. The laser saw technique is an alternative to be used to cut the epitaxy layer into the regular shape. In application of the laser chemical processing method, the cutting loss may be FIG. 2.5.9 Schematic diagram of the setup for laser polishing [34].

2.5.3 Single-crystal diamond wafer process [4, 8]

193

FIG. 2.5.10 Shows the procedure of lift off [35].

reduced by using a water jet-guided laser beam. A very thin liquid jet stream (width 30–80 micrometers) is generated in a nozzle that has a transparent window on top. The laser beam is focused into the jet stream and guided via total internal reflection. The material removal occurs by heating or evaporation and can be enhanced by adding etching solutions. This also offers the possibility to remove any surface damage directly. As diamond wafer relies on growth, the lift-off process becomes the main process for diamond wafer to be separated from the substrate instead of the conventional cutting process. The lift-off process is based on ion implantation, which consists of three steps: (1) carbon ion implantation to form a buried, well-defined damaged layer below the surface, (2) graphitization of this layer by annealing in vacuum, and (3) etching of the damaged layer in an acid solution. By an improved process, large (10  10 mm) and thick (>0.2 mm) single-crystal diamond plates were lifted off, and the substrate could be used as a seed for more than three times without polishing the surface. Fig. 2.5.10 shows the procedure of lift off [35].

2.5.3.3 Grinding and lapping [12] In the conventional wafer routine, the as-cut surface of the ingot always consists of subsurface damage such as microcracks, dislocations, and deformed or phase-transformed layers. Therefore, a grinding and/or lapping procedure is performed on these as-cut surfaces of the ingot to remove these subsurface damages. After this (or these) procedure(s), the surface of the treated slice is expected to obtain the initial global planarization and the required surface qualities. The surface conditions are usually specified by roughness, texture, and planarity. The obvious difference between grinding and lapping is the use of an abrasive. In the lapping procedure, abrasives are suspended in slurry to perform the mechanical process of removing the surface roughness and subsurface damage. So, the lapping procedure is based on the mechanical free abrasive machining (FAM) process while the fixed abrasive grinding wheel is usually used in the grinding procedure. It has high throughput and fully automatic operation over the lapping procedure. However, the spring-back effect in grinding is much larger than that in lapping. This may result in no homogenous thickness of the wafer. Therefore, lapping has to be followed after grinding. The relationship between grinding and lapping is schematically seen in Fig. 2.5.11 [5]. As a large-size single-crystal diamond is still difficult to support, lapping is much more popular than grinding.

194

2. Semiconductor diamond

Grinding Sliced wafer

Polishing Lapping

Surface grinding

Etching

Two-sided grinding Prepolishing Soft-pad grinding

Other grinding processes Ground wafers

Polishing

Polished wafers

FIG. 2.5.11 The relationship between grinding and lapping [5].

2.5.3.4 Polishing Polishing is the procedure to perform the macroflattening and microsmoothing of the treated surfaces after grinding and/or lapping so that the surfaces can arrive at the required wafer surfaces. In this procedure, various methods are used including mechanical, chemical, electrolytic, or thermal methods or even the synergistic combination of these methods to increase the polishing efficiency and reduce the economic cost. An appropriate polishing technique must be selected based on the specific application requirement, the shape of the workpiece, and the existing equipment. The material removal rate and the roughness of the polished surface is always used to evaluate the polishing technique. There are several proposed methods for single-crystal diamond polishing, including mechanical polishing (MP), chemomechanical polishing (CMP), thermochemical polishing (TCP), high energy beam (laser/plasma/ion beam) polishing, dynamic friction polishing (DFP), electrical discharge machining, and several other polishing techniques. Among them, CMP, TCP, and LP have been introduced for polycrystalline grain removal. Therefore, MP, DFP and IBP will be the focus. 2.5.3.4.1 Mechanical polishing Mechanical polishing (MP) is a contact polishing method that is commonly used for polishing single-crystal diamond. The efficiency of MP is affected very much by the crystallographic orientation of the diamond surface being polished and the direction of sliding because of the polishing resistance. The easily polished crystallographic planes are the (100) and (110) planes, but it is extremely difficult to polish the (111) plane. Even on the easily polished crystallographic planes, the polishing should be performed along the soft direction or easy

2.5.3 Single-crystal diamond wafer process [4, 8]

195

direction, avoiding the hard direction. The soft direction refers to polishing in the <100> or < 110 > direction on the (100) and (110) planes, and all directions on the (111) surface are termed hard direction. The material removal rate is influenced by the relative sliding speed and load (hence of pressure) as well as the size and concentration of diamond grit on the scaife and the environmental conditions. Generally speaking, the high rotation speed of the scaife and the large load exerted on the polished surface can increase the material removal rate. It shows the linear relationship at medium speed as about 100m/s; the nonlinear relationship occurs at low speed and very high speed because of the property change in the iron-casted scaife. The reason for the load affecting the material removal rate may lie in the friction coefficient change in a different direction with the load. The larger diamond powder can obtain a higher material removal rate while a finer powder results in a higher quality finish. In regard to concentration, recharging made no significant difference in an easy < 100 > direction. But in the hard direction, frequent recharging greatly increased the polish rate. Mechanical polishing is highly accurate and commonly used, although the polishing rate is not fast and sometimes it is impossible to polish in hard directions. Although some certain chemical reactions are utilized to increase the material removal rate, the material removal rate is not obviously improved. For example, the oxidation reaction is utilized by coating the amorphous silicon oxide (SiOx) on a rotating scaife. The SiOx (x¼ 1.97) reacts with the diamond to form CO and CO2 when the diamond is rubbed on the scaife. The diamond surface is thus polished due to the carbon being removed from the surface chemically and mechanically. Polishing a single-crystal diamond requires great skill because the soft direction on the easily polished crystallographic planes needs to examine by the experience. By MP, the surface roughness can arrive at several angstroms, and the material removal rate can arrive at several tens of μm/h. The typical equipment is shown in Fig. 2.5.12. 2.5.3.4.2 Dynamic friction polishing Like MP, another contact polishing technique is dynamic friction polishing (DFP), which is a straightforward technique for the polishing of single-crystal diamond. DFP is an abrasiveFIG. 2.5.12

The instrument of mechanical polishing for single-crystal diamond jewelry.

196

2. Semiconductor diamond

free process that utilizes the frictional heating between a diamond specimen and a rotating catalytic metal disk. As the thermochemical reaction is activated, the conversion of diamond into nondiamond carbon, diffusion, oxidation, and mechanical abrasion take place to achieve an efficient material removal. For example, a quality surface finish of the roughness of 50 nm Ra can be obtained in some minutes. The equipment required is simple, and the process is relatively easier to control in a normal ambient environment without a vacuum chamber and/or special heating. So, it is a cost-effective technique. Polishing parameters are the disk material, the pressure being exerted on the diamond, the rotating speed/sliding speed, and the polishing environments. Among these parameters, the pressure pressing on the diamond is a key factor, that is, there is a threshold value of pressure to start the polishing at a certain rotating speed/sliding speed. If the pressure is lower than the threshold value, DFP won’t start. Blowing oxygen or air into the polishing point will enhance the polishing rate if polishing is starting. The polishing rate is 100 times faster than the other polishing techniques, and the polished surface possesses a roughness of 50 nm of Ra. In summary, the significant features of the DFP method and its potential superiority over the other existing methods include: (1) Abrasive-free polishing by simply pressing a diamond workpiece to a metal disk rotating at a high speed to generate frictional heat. (2) Polishing at room temperature, effective use of polishing-generated frictional heating, and no requirement for special equipment to heat the polishing disk and the diamond specimen. (3) Polishing in a dry atmospheric environment and no requirement for a vacuum chamber. (4) Short polishing time (on the order of a few tens of seconds to a few minutes), material removal rate is up to 480 μm/min for single-crystal diamond. (5) Surface roughness values of the order of 50 nm Ra for single-crystal diamond. The schematic diagram of the setup for DFP is shown in Fig. 2.5.13 [36]. FIG. 2.5.13 Schematic diagram of the setup for dynamic friction polishing [36].

Load Metal disk

Specimen holder

PCD specimen

197

2.5.3 Single-crystal diamond wafer process [4, 8]

2.5.3.4.3 Ion beam polishing Ion/plasma beam polishing (IBP/PBP) is a noncontact polishing technique, also called high energy beam polishing, which applies the principle of bombardment of the diamond surface with reactive/nonreactive ion beams to ablate diamond atoms by sputtering. Two courses result in the removal of the carbon atoms. First, the heating due to the bombardment of the incident ion sublimates the host lattice atoms. Second, the incident ion can simply knock off a host atom by sputtering. To enhance the material removal, oxidation and etching are also utilized, as in reactive ion etching (RIE), that is, the reactive channels wherein O, O2, O3, OH, H, etc., interact and remove carbon from the diamond surface via heterogeneous chemical reactions. In ion beam polishing, polishing rates and surface finish are dependent on the incidence angle of ion irradiation with respect to the diamond surface, the energy power of irradiating, the type of ion beam and gap flow, and the chamber vacuum pressure. As the polishing rate is not very fast, IBP/PBP can be regarded as the polishing method for a surface fine finish of the single-crystal diamond wafer instead of a coarse polycrystalline diamond film. It means that IBP/PBP is not suitable for polycrystalline grain removal in the production of single-crystal diamond wafer. The instrument is shown in Fig. 2.5.14 [37]. 2.5.3.4.4 Other polishing techniques There are some other methods for polishing diamond, such as electrical discharge machining (EDM), abrasive liquid jet polishing, aero-lap polishing, float polishing, spark erosion, etc. However, these methods are not suitable for single-crystal diamond wafer polishing. For example, the EDM polishing method is suitable for electrical conductive PCD instead of nonelectrical conductive single-crystalline diamond and CVD polycrystalline diamond films. 2.5.3.4.5 Comparison and selection of polishing techniques (1) Comparison of polishing techniques

Beam Shutter Work-stage

Beam extraction electrode

RF-Generator 2Mz

Magnet

Einzel lenz Ions

Working chamber (10–4 Pa)

DC Discharge chamber

Gas Metal barrier

Turbo-molecular pump DC-Generator

(A)

(B)

FIG. 2.5.14 (A) An ion beam machining apparatus equipped with a high-voltage discharge-type ion source [37]. (B) Schematic diagram of a plasma polishing system [38].

198

2. Semiconductor diamond

Table 2.5.3 summarizes the features and characteristics of the various polishing techniques [8]. A polishing technique should be selected based on the quantity of the diamond sample to be polished, the final surface finish required, the polishing process time, and the equipment cost. Optimization of polishing parameters is often essential to ensure the best possible results. MP is a relatively straightforward and scalable process, and there is no requirement for substrate heating or using reactive gases. It is widely used in industry for polishing single-crystal diamond on selected soft directions. This method produces a polished surface with average surface roughness of the order of a few nanometers, and the polishing does not drastically change the chemical quality of the diamond surface. However, this polishing method has extremely low polishing rates that depend on the quality of the diamond and its lattice orientation. Sometimes, the polishing time can be quite long, even up to several days. In addition, the applied pressure on the substrates can cause microcracking on a thin CVD single-crystal diamond wafer, and the consuming abrasive diamond or diamond wheel is expensive. CMP is suitable for polycrystalline grain removal, and it also can provide a higher material removal rate, better surface roughness, and less surface damage than mechanical polishing in single-crystal polishing. However, the heating of the polishing disk and the requirement of oxidizing agents make the polishing process complicated. In addition, to maintain a continuous polishing process, it is essential to remove the chemically reacted products accumulated on the polishing disk. TCP is also used to remove polycrystalline grains. It can also offer a fine surface finish in single-crystal diamond polishing. The advantages of TCP are that the ultrafine surface could be obtained with a roughness of about 2 nm without encountering difficulties in defined planes and directions while the polishing rate is much higher than that of MP. The disadvantage is that the equipment and the procedure are complex because efficient polishing can only be achieved by heating the polishing disk to temperatures higher than 750°C, which requires an evacuated/reductive atmosphere to prevent the metal from being oxidized, especially when using iron at high temperatures. Thus, other problems relating to the facilities arise that complicate the polishing process. In addition, surface nonuniformities are introduced partially from the mating metal surface, and contamination occurs with the formation of a diamond-like carbon layer and metal residues in the grain boundaries. The significant features of the laser-based polishing method and its potential superiority over the other existing methods include: polishing at room temperature and a short polishing time (from a few seconds to a few minutes); it is a noncontact process and there is no restriction on the shape of the surface to be polished; and a very fine surface finish can be achieved. Laser polishing has advantages that can be attractive for particular applications: (a) localized polishing of selected small areas; (b) polishing of larger areas inside any well-defined boundary by multipulse beam scanning; and (c) polishing of curved surfaces. On the other hand, laser polishing has the following disadvantages: chemical nonuniformity of the surface limits the polishing process; the polishing area is restricted to laser spot size and scanners are required for polishing large areas of contoured surfaces; and graphitic layers could contaminate the surface after polishing of the diamond. Ion beam polishing methods can achieve a very fine surface with a roughness value of the order of a few nanometers or less; they are very suitable for many electronic device applications requiring precision polishing in small areas. The polished surface is clean and does not

TABLE 2.5.3 Polishing technique

The features and characteristics of the various polishing techniques [8] Mechanical

Chemomechanical

Hot metal plate

Thermoetching

Dynamic friction

Ion beam Laser

laser

EDM

Contact

Reactive contact

Reactive contact

Reactive contact

Reactive contact

Noncontact

Noncontact

Noncontact

Bulk processing temperature

Room

Generally over 300 °C

750–950°C

700–950°C

Room, Friction heat

Room or 700°C for RIE

Room

Room

Polishing mechanisms

Abrasive wear

Abrasive wear, oxidation

Graphitization diffusion

Diffusion

Graphitization Mechanical wear chemical reaction

Sputtering

Evaporation

Evaporation

Special requirement

None

Oxidation protection

Vacuum or reduction gas

Vacuum or reduction gas

None

High vacuum Gas flow

Scanning of the sample

None

Set-up

Rigid and geometry sensitive

Complex

Heating chamber

Medium

Rigid and geometry sensitive

Complex

Simple

Geometry sensitive

Equipment cost

Low, but consuming diamond abrasives

Medium

Medium

Low

Low

High

High

Medium

Large area processing cost

Medium

Medium

Medium

Low

Low

High

Medium

Low

Material and surface finish

Precise for singlecrystal diamond (A few nm Ra)

Ultrap recise for PCD, Single, and CVD DF (< 1 nm Ra)

Ultrafine on CVD DF(A few nm) Ra

Rough for CVDDF (>1000 nm)

Precise for single diamond (50 nm Ra) Good for PCD (200 nm)

Precise for PCD, Single and CVD DF (a few nm Ra)

Good for CVD DF and PCD (5–70 nm Ra)

Rough for PCD and CVD DF (>500 nm Ra)

Polishing rate

Tens of um/h

Few μm/h

Few μm/h

Tens of μm/h Simultaneous for many

1000 μm/h

Tens of μm/h

Hundreds of μm/h

Few μm/ min

199

Continued

2.5.3 Single-crystal diamond wafer process [4, 8]

Nature of processing

The features and characteristics of the various polishing techniques [8]—cont’d

200

TABLE 2.5.3

Mechanical

Chemomechanical

Hot metal plate

Thermoetching

Dynamic friction

Ion beam Laser

laser

EDM

Processing time (per cm2)

Tens of hours

A few hours

Few hours

Few hours

Few minutes

Tens of hours

A few to tens seconds

Few minutes

Shape limitation

Planar surfaces

Planar surfaces

Planar surfaces

Planar Nonplanar

Planar surfaces

Planar Nonplanar

Nonplanar

Nonplanar

Size limitation

Disk size

Plate size

Plate size

No limit

Disk size

Beam size

No limit

No limit

Surface change orcontamination

Little

Medium

Medium

High

Medium

Little

Medium

Medium

Potential for commercialization

High

Medium

Medium

Medium

High

Medium

High

High

Anisotropic feature

High

Low

Low

Medium

Medium

Low

–

–

2. Semiconductor diamond

Polishing technique

References

201

leave chemical contamination on diamond. However, this technique suffers from a complicated experimental setup and large capital cost coupled with critical requirements for orientation of the sample with respect to ion beam incidence. This restricts the area of polishing to the size of the ion beam. Also, these methods achieve highly nonuniform polishing and contamination in the form of residue formation on the surface between grain boundaries. (2) Selection of polishing techniques Single-crystal diamond can be readily polished by traditional mechanical polishing to a surface roughness of the order of a few nanometers. As the polishing is highly anisotropic, the polishing rate would be much higher in a soft direction on a given plan. Polishing a single-crystal diamond requires great skill as polishing is carried out while examining the crystallographic planes and orientation to locate the plane to be straightforwardly polished. In addition, polishing of single-crystalline diamond by using the DFP method can dramatically increase the polishing efficiency. The surface roughness can be down to the order of 50 nm in a few minutes. If a better surface finish is required for some application, the diamond surface should be further polished by a mechanical polishing process using small diamond abrasives, or by using a chemomechanical method to obtain an ultrafine surface.

2.5.4 Conclusion (1) To obtain a large-sized wafer, the key procedure is to perform large crystal growth because wafer processing is the “reducing volume” procedure after crystal growth. Taking into account the high cost and hardness of the row material, diamond wafer may consider the net-shape manufacturing. In this procedure, the diamond wafer can only be grown without wafer processing, or with part of wafer processing. It may be a possible approach to obtain a large-sized single-crystal diamond wafer. (2) To obtain wafer from CVD diamond, the first procedure is to remove the polycrystalline grains from the as-grown surface. In this procedure, it is suitable to consider the application of CMP, TCM, and even LP. (3) In the grinding/lapping procedure, lapping is much more popular than grinding due to not having a large enough slice of row diamond. (4) MP, DFP, and IBP/PBP are suitable for single-crystal diamond wafer polishing. The efficiency is the highest for DFP, and IBP/PBP is mild for the polishing surface. MP may result in much more surface damage, so the following treatment should be added to remove the spoiled layer.

References [1] S. Matsumoto, Y. Sato, N. Setaka, Effect of the preceding heat treatment on hydrogen chemisorption of diamond powders, Carbon 19 (3) (1981) 232. [2] C.S. Yan, Y.K. Vohra, H.K. Mao, R.J. Hemley, Very high growth rate chemical vapor deposition of single-crystal diamond, Proc. Natl. Acad. Sci. U. S. A. 99 (20) (2002) 12523. [3] S.T. Lee, Y. Lifshitz, The road to diamond wafers, Nature 424 (6948) (2003) 500. [4] H.–.J. M€ oller, P. Rudolph (Eds.), Wafer processing, in: Handbook of Crystal Growth, Elsevier Inc., 2015.

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[5] I. Kao, C. Chung, R.M. Rodriguez, Wafer Manufacturing and Slicing Using Wiresaw. in: G. Dhanaraj, K. Byrappa, V. Prasad, M. Dudley (Eds.), Springer Handbook of Crystal Growth, Springer, Heidelberg, Dordrecht, London, New York, Berlin, ISBN: 978-3-540-74182-4, 2010. e-ISBN: 978-3-540-74761-1. https://dx.doi.org/10. 1007/978-3-540-74761-1. [6] T.N. Bhat, S.B. Dolmanan, Y. Dikme, H.R. Tan, L.K. Bera, S. Tripathy, Structural and optical properties of AlxGa1-xN/GaN high electron mobility transistor structures grown on 200mm diameter Si(111) substrates, J. Vac. Sci. Technol. B 32 (2) (2014) 21206. [7] http://www.nssmc.com/en/news/old_nsc/detail/index.html/?rec_id¼4177. [8] Y.Q. Chen, L.C. Zhang, Polishing of Diamond Materials Mechanisms. in: B. Derby (Ed.), Modeling and Implementation, Engineering Materials and Processes1619-0181Springer, London, Heidelberg, New York, Dordrecht, 2013. ISBN 978-1-84996-407-4; ISBN: 978-1-84996-408-1 (eBook). Library of Congress Control Number: 2012955039. https://dx.doi.org/10.1007/978-1-84996-408-1. [9] L. Pastewka, S. Moser, P. Gumbsch, M. Moseler, Anisotropic mechanical amorphization drives wear in diamond, Nat. Mater. 10 (1) (2011) 34. [10] H.D. Li, Epitaxy of carbon-based materials: diamond thin film, Chapter 14, in: T.F. Kuech (Ed.), Handbook of Crystal Growth: Thin Films and Epitaxy, Elsevier Inc, 2015. ISBN: 978-0-444-63304-0. [11] Y.N. Palyanov, I.N. Kupriyanov, A.F. Khokhryakov, V.G. Ralchenko, Crystal growth of diamond, Chapter 17, in: T.F. Kuech (Ed.), Handbook of Crystal Growth: Thin Films and Epitaxy, Elsevier Inc, 2015. ISBN: 978-0-44463304-0. [12] O. Ohnishi, H. Suzuki, E. Uhlmann, N. Schr€ oer, C. Sammler, G. Spur, M. Weismiller, Grinding, Chapter 4in: I. D. Marinescu, T.K. Doi, E. Uhlmann (Eds.), Handbook of Ceramics Grinding and Polishing, Elsevier Inc., 2015. ISBN: 978-1-4557-7858-4. [13] S. Ohmagari, H. Yamada, H. Umezawa, A. Chayahara, T. Teraji, S.-i. Shikata, Characterization of free-standing single-crystal diamond prepared by hot-filament chemical vapor deposition, Diam. Relat. Mater. 48 (2014) 19. [14] G.C. Chen, B. Li, H. Li, H. Lan, F.W. Dai, Q.J. Xue, X.Q. Han, L.F. Hei, J.H. Song, C.M. Li, W.Z. Tang, F.X. Lu, Growth of diamond by DC Arcjet Plasma CVD: from nano-sized poly-crystal films to millimeter-sized single crystal grain, Diam. Relat. Mater. 19 (2010) 1078. [15] M. Schreck, S. Gsell, R. Brescia, M. Fischer, Ion bombardment induced buried lateral growth: the key mechanism for the synthesis of single-crystal diamond wafers, Sci. Rep. 7 (44462) (2017) 1. [16] J. Liu, L.F. Hei, J.H. Song, C.M. Li, W.Z. Tang, G.C. Chen, F.X. Lu, High-rate homoepitaxial growth of CVD single-crystal diamond by dc arc plasma jet at blow-down (open cycle) mode, Diam. Relat. Mater. 46 (2014) 42. [17] S.T. Lee, Z.D. Lin, X. Jiang, CVD diamond films: nucleation and growth, Mater. Sci. Eng. R Rep. 25 (4) (1999) 123. [18] H. Yamada, A. Chayahara, Y. Mokuno, S.–.I. Shikata, Diam. Relat. Mater. 17 (7–10) (2008) 1062. [19] C. Jany, A. Tardieu, A. Gicquel, P. Bergonzo, F. Foulon, Diam. Relat. Mater. 9 (2000) 1086. [20] C.S. Yan, Y.K. Vohra, Diam. Relat. Mater. 8 (1999) 2022. [21] Y. Mokuno, A. Chayahara, Y. Soda, Y. Horino, N. Fujimori, Diam. Relat. Mater. 14 (11–12) (2005) 1743. [22] K. Kurihara, K. Sasaki, M. Kawarada, N. Koshino, High rate synthesis of diamond by dc plasama jet chemical vapor deposition, Appl. Phys. Lett. 52 (1988) 437. [23] G.C. Chen, B. Li, Z.Q. Yan, F.X. Lu, Single crystalline diamond grown by repositioning substrate in DC arcjet plasma enhanced chemical vapor deposition, Diam. Relat. Mater. 21 (2012) 83. [24] N. Ohtake, M. Yoshikawa, Diamond film preparation by arc discharge plasma jet chemical vapor deposition in the methane atmosphere, J. Electrochem. Soc. 137 (1990) 717. [25] H.L. Michael, M.A. Cappelli, Diamond synthesis in supersonic direct-current arcjet plasma at subtorr pressures, Surf. Coat. Technol. 54–55 (2) (1992) 408. [26] F.X. Lu, G.F. Zhong, J.G. Sun, Y.L. Fu, W.Z. Tang, J.J. Wang, G.H. Li, J.M. Zang, C.H. Pan, C.X. Tang, T.L. Lo, Y. G. Zhang, A new type of DC arc plasma torch for low cost large area diamond deposition, Diam. Relat. Mater. 7 (6) (1998) 737. [27] L.F. Hei, J. Liu, C.M. Li, J.H. Song, W.Z. Tang, F.X. Lu, Fabrication and characterizations of large homoepitaxial single-crystal diamond grown by DC arc plasma jet CVD, Diam. Relat. Mater. 30 (2012) 77. [28] F.X. Lu, W.Z. Tang, T.B. Huang, J.M. Liu, J.H. Song, W.X. Yu, Y.M. Tong, Diam. Relat. Mater. 10 (2001) 1551. [29] H. Yamada, A. Chayahara, H. Umezawa, N. Tsubouchi, Y. Mokuno, S. Shikata, Fabrication and fundamental characterizations of tiled clones of single-crystal diamond with 1-inch size, Diam. Relat. Mater. 24 (2012) 29. [30] S. Matsumoto, M. Hino, T. Kobayashi, Synthesis of diamond films in a rf induction thermal plasma, Appl. Phys. Lett. 51 (1987) 737.

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[31] J.–.J. Li, B. Li, Y.–.G. Zuo, H. Liu, Y. Bai, H.–.W. Yuan, Z.–.R. Li, K. Xua, G.C. Chen, Application of dual radio frequency inductive coupled plasma into CVD diamond growth, Vacuum 154 (2018) 174. [32] T. Doi, I.D. Marinescu, S. Kurokawa, Chapter 4—Applications of ultra-precision CMP in device processing, in: Advances in CMP Polishing Technologies, Elsevier Inc, 2012, pp. 113–228. [33] H.Y. Tsai, C.J. Ting, C.P. Chou, Evaluation research of polishing methods for large area diamond films produced by chemical vapor deposition, Diam. Relat. Mater. 16 (2) (2007) 253. [34] C. Chen, H.–.L. Tsai, Fundamental study of the bulge structure generated in laser polishing process, Opt. Lasers Eng. 107 (2018) 54. [35] Y. Mokuno, A. Chayahara, H. Yamada, N. Tsubouchi, Improvements of crystallinity of single-crystal diamond plates produced by lift-off process using ion implantation, Diam. Relat. Mater. 19 (2–3) (2010) 128. [36] Y. Chen, L.C. Zhang, Polishing of polycrystalline diamond by the technique of dynamic friction, part 4: Establishing the polishing map, Int. J. Mach. Tools Manuf. 49 (3–4) (2009) 309. [37] Y. Sato, J. Kawamura, T. Nagase, S.A. Pahlovy, I. Miyamoto, Sharpening of CVD diamond coated tools by 0.5 10 keV Ar+ ion beam, Diam. Relat. Mater. 20 (7) (2011) 954. [38] E.E. Yunata, T. Aizawa, K. Tamaoki, M. Kasugi, Plasma polishing and finishing of CVD-diamond coated WC (Co) dies for dry stamping, Procedia Eng. 207 (2017) 2197.

C H A P T E R

2.6 Impurities and dopants in diamond Meiyong Liao Research Center for Functional Materials, National Institute for Materials Science (NIMS), Tsukuba, Japan

Intrinsic diamond is a perfect insulator at room temperature due to the ultrawide bandgap energy. To enable diamond as a semiconductor, it is necessary to dope it with foreign atoms. Because the lattice constant of diamond is small, the foreign atoms able to incorporate into the bulk to form electrically active dopants are limited. Natural diamonds contain a large number of impurities such as metals, hydrogen, oxygen, boron, nitrogen, etc., most of which are not electrically active [1]. The most common impurity in artificial and natural diamonds is nitrogen. A small amount of diamond (0.1%) with boron impurity also exists in nature [2]. According to the amount of nitrogen and the arrangement in the lattice, natural diamonds can be classified into type I (nitrogen) and II (less nitrogen), specifically [3], as show in Fig. 2.6.1 [4]. Type Ia: More than 95% of natural diamonds belong to this type, which contains a large of amount of nitrogen with a concentration up to 3000 ppm (0.3%). The nitrogen within a type-Ia diamond is aggregated into two forms: A-aggregates (IaA), which consist of pairs of nitrogen

204

FIG. 2.6.1

2. Semiconductor diamond

Types of diamond [4].

atoms, and B-aggregates (IaB), which are composed of four N atoms around a vacancy (V). Optically, type-Ia diamonds often show sharp absorption bands of N2 (478 nm) and N3 (415.5 nm) centers accompanied by weaker lines at 465, 452, 435, and 423 nm related to N2 and N3 centers. Type-Ia diamonds show blue to UV fluorescence due to the N3 centers. There are also H3 centers at 504 nm, sometimes accompanied by an H4 center at 537 and 495 nm. Type-Ia diamonds exhibit various colors such as colorless, brown, pink, and violet. The color can be changed to yellow, orange, red, blue, and green by radiation or thermal treatments [5, 6]. Type Ib: In these diamonds, nitrogen atoms replace carbon atoms in the lattice and are isolated from one another. One cannot find nitrogen atoms in adjacent lattice positions. The nitrogen concentration is up to 500 ppm (0.05%). Type-Ib diamonds have a yellow color. These nitrogens are called single substitutional nitrogen or isolated nitrogen. The substitution of nitrogen for carbon produces an unpaired electron localized between nitrogen and carbon, which shows a paramagnetic resonance effect. About 0.1% of natural diamonds and most of the synthetic diamonds are this type. Type IIa: This type of diamond has little nitrogen and is colorless. Around 1% of natural diamonds belong to this type. Type IIb: Natural type-IIb diamonds contain little nitrogen (not measured by infrared light). Instead, a large amount of boron impurities, which replace carbon atoms in the lattices, appears in this type of diamond. Type IIb diamonds show high p-type electrical conductivity. Nitrogen can easily be doped as a donor in diamond, which, however, shows a very deep activation of 1.7 eV, even larger than the bandgap of Si. Although p-type diamonds exist in nature, the electrical conductivity and crystal quality cannot be controlled. In bulk diamond, boron is realized as the only acceptor and phosphorous the only donor up to now. In addition, hydrogenated-terminated diamond exhibits high p-type surface conductivity. Nowadays, ptype diamond can be readily achieved through either bulk or surface doping while n-type diamond can only be achieved in limited conditions. Unfortunately, both acceptor boron (0.37 eV) and donor phosphorous (0.57 eV) show high thermal activation energies in diamond. Surface p-type conductivity is the most explored for semiconductor electronic devices due to the high sheet density of 1013 cm2.

2.6.1 Bulk p-type diamond

205

2.6.1 Bulk p-type diamond Due to the deep level of boron in diamond, only a fraction of the boron atoms can contribute to the electrical conductivity at room temperature. Full activation of boron can only be achieved at very low boron concentration when the Fermi level crosses the acceptor level or at high concentrations, when the miniband is going to the valence band, as shown in Fig. 2.6.2A [8]. The full activation of boron occurs when the boron concentration is above 1020 cm3. A resistivity around 1 mΩ cm was achieved at 300 K when the boron concentration was above 1021 cm3 [9]. For a nondegenerate semiconductor where one acceptor and one compensating donor are present with concentrations Na and Nd, respectively, and with Na > Nd, the hole concentration P can be calculated as the following:   pðP + Nd Þ  n2i Nv EA ¼ exp  KT Na  Nd  p  n2i ga where NV is the effective density of the states of the valence band, ni is the intrinsic carrier concentration, ga is the spin degeneracy factor, k is the Boltzmann constant, Ea is the acceptor ionization energy, and T is the absolute temperature. The hole concentrations at room temperature for different boron doping concentrations and different compensating donor concentrations are shown in Fig. 2.6.2B [7]. Bulk p-type diamond is normally achieved by adding the boron source during the CVD growth. The boron source used for a boron-doped p-type diamond was B2H6 at the early stage [10]. The hole mobility was reported to be around 1370 cm2 V1 s1 for a boron concentration close to 1017 cm3 and larger than 1000 cm2 V1 s1 for a boron concentration of 1017 cm3. Nowadays, less toxic gas [B(CH3)3, TMB] diluted in H2 has been widely used for bulk p-type diamond doping [11, 12]. The hole mobility achieved by TMB was 1800 cm2V1 s1 at 290 K for a boron concentration of 6 1017 cm3 (hole concentration 1014 cm3), comparable to that achieved by using B2H6 as the boron source. By adding oxygen in the gas during the CVD growth, the boron concentration in the diamond epilayer can be reduced to less than 1015 cm3 [13]. The mobility in a low boron concentration close to 1016 cm3 showed a Hall mobility of 1870 cm2/V s at 292 K [14]. Due to the deep nature of boron in diamond, the hole concentration in boron-doped diamond is varied as temperature in the medium boron concentrations. As shown in Fig. 2.6.3, the hole concentration changes exponentially with reciprocal temperature. One can estimate the activation energy of the boron acceptor in diamond, which is around 0.37 eV between 170 and 270 K [12]. Hole mobility is also a function of temperature, which follows the law of T1.9 with temperature above 1017 cm3 220 K, as illustrated in Fig. 2.6.4. In a different study on the Hall mobility dependence on temperature for a [B] close to 1016 cm3, a similar trend was also observed. Because the boron concentration was low, ionized impurity scattering was not considered as the dominated mechanism. Instead, the crystal quality limited the hole mobility [14]. Doping of boron in diamond also generates dislocation defects [15]. Depending on the diamond facet, there are critical boron concentrations for defect generation. For the (100) diamond, the critical boron centration for defect generation is 3.2  1021 cm3 and for the (111) diamond, the critical value is 6.5–17 1020 cm3.

206

2. Semiconductor diamond

FIG. 2.6.2 (A) Activation ratio of the boron concentration for different temperatures and (B) hole concentration with different boron concentrations and donor compensations [7].

Implantation of boron ions inside diamond followed by high-temperature annealing provides an alternative method for p-type conductivity [16]. In this method, low-dose (1014 cm2) ions have to be adopted in order to avoid the graphitization. A high temperature annealing as high as 1400oC was also required to recover the lattice damage. Prawer showed a boron concentration of 1018 cm3 and a hole mobility of 585 cm2/V s by 2MeV B ion implantation with a dose of 1015 cm2 at a low temperature of 77 K and annealed at 1450oC [17]. The MeV implantation can also be conducted at room temperature followed by a high temperature annealing [18].A fraction of 15%–30% of the implanted boron could be activated at high temperatures for an implantation dose of 1015 and 1014 cm2, respectively. The hole mobility was reported to be about 230 cm2/V s after annealing at 1500°C.

2.6.1 Bulk p-type diamond

FIG. 2.6.3

Hole concentration as temperature for a borondoped diamond with a concentration [B]¼61017 cm3 [12].

Temperature (K) 300

1016

250

207

200

Hole concentration (cm–3)

1015 1014 1013 1012 1011 1010 109 2.5

3

3.5

4

4.5

5

5.5

6

1000/T (K–1)

FIG. 2.6.4 4000 Hall mobility (cm2/Vs)

Dependence of hole mobility as temperature

[12].

3000

2000

1000 100

300 200 Temperatrue (K)

400

Recently, a novel method based on thermal diffusion was reported [19]. In this method, substitutional doping of boron in SCD was achieved through annealing a boron-doped Si nanomembrane bonding on an SCD plate (Fig. 2.6.5). The boron was distributed with 50 nm of the SCD surface. A boron concentration of 2 1018 cm3 and a hole mobility of 120 cm2/V s were achieved. The authors claimed that no graphitization was observed. Schottky diodes with a breakdown voltage of 80 V were demonstrated by using the thermal-diffusion doping method.

208

2. Semiconductor diamond

FIG. 2.6.5

(A) Process for boron doping of SCD by the thermal diffusion process. (i) Formation of heavily borondoped Si on the SOI substrate by boron implantation and thermal annealing. (ii) Etching of SiO2 to form silicon nanomembranes (SiNM). (iii) Picking up the top Si by an elastomeric stamp. (iv) Transfer of the SiNM to a diamond plate. (v) Rapid thermal annealing (RTA) of the SiNM on diamond to facilitate the thermal diffusion with RTA. (vi) SiNM removed by potassium hydroxide (KOH) etching. (B) Microscopic image of diamond plate before SiNM transfer. The inset shows a zoomed-in image of part of the diamond plate. (C) Image of diamond plate bonded with SiNM strips [19].

In order to achieve sufficient electrical conductivity for device application, heavily borondoped p-type diamond is necessary [20]. However, a high boron connotation degrades the hole mobility markedly. δ-doping with a profile of within 2 nm buried in the intrinsic diamond was proposed to achieve both high conductivity and high mobility [21]. The growth of δ-doped diamond with a sharp profile is quite challenging, which requires the optimization of the growth conditions from the substrate to the CVD chamber design. To achieve an ideal δ-doping, one needs to grow a very thin (1–2 nm), heavily doped δ layer. In addition, the interfaces between the δ layer and the high mobility intrinsic or lightly boron-doped diamond layers must be atomically smooth to minimize carrier scattering. Most of the efforts in δdoped diamond showed a low hole mobility of less than 10 cm2/V s when boron concentration was up to 5 1020 cm3 [22]. Recently, Bulter et al. achieved a sharp δ  doping p-type layer with a Hall mobilities of up to 120 cm2/V s and a sheet carrier concentration up to 6  1013 cm2. The δ-doped profile had a profile with a 2-nm FWHM. These are the best values up to now. Such a breakthrough in δ-doped diamond relied on: [23] (a) the use of a rapid gas switching system with a time of less than 10 s, (b) the reactor design providing the laminar gas flow, (c) very slow growth rates, (d) smooth epitaxial films, and (e) chemically gettering residual boron in the reactor when growing intrinsic/lightly doped material by adding H2S to form gaseous boron sulfur compounds. Both the single δ  doped layer and multiple δ  doped layers were obtained.

209

2.6.2 p-Type surface conductivity and transfer doping

2.6.2 p-Type surface conductivity and transfer doping A unique property of diamond that differs from other conventional semiconductors is the appearance of p-type surface conductivity with a sheet hole concentration of 1013 cm2 and a hole mobility of 100 cm2/V s when an intrinsic diamond with hydrogen (H) termination is exposed to air [24, 25]. The surface conductivity has been realized to have the twodimensional hole gas (2DHG) nature [26]. The hole conductivity maintains even at a very low temperature of 0.34 K [27]. The H-termination of diamond is formed during the diamond growth by CVD, where H2 gas is normally fed. The H-terminated diamond surface exhibits a negative electron affinity (NEA) of 1.3 eV, which lowers the ionization to 4.2 eV [28]. When putting an appropriate adsorbate with high electron affinity on the H-diamond surface, the electron can transfer from the valence band of diamond into the lowest unoccupied molecular orbital (LUMO) of the absorbate, as shown in Fig. 2.6.6. Such a surface transfer enables the formation of a hole accumulation layer on the diamond surface [30, 31]. Later, it was shown that p-type surface doping could also be achieved with fullerene (C60) and fluorinated fullerenes C60F48 [32, 33]. The surface conductivity resulting from H-diamond exposure to air or molecular has been widely adopted for the fabrication of diamond MOSFETs [34, 35]. However, the surface conductivity of H-terminated diamond due to ambient air exposure suffers from instability upon environmental conditions such as temperature, humidity, and molecular composition [36]. Passivation of the H-terminated diamond surface with Al2O3 greatly enhanced the thermal stability of p-type conductivity up to 700 K [37]. The deposition of high electron affinity oxides such as Nb2O5, WO3, V2O5, and MoO3 can also generate p-type surface conductivity on H-terminated diamond through surface transfer [38–41]. The sheet density is comparable to that of H-terminated diamond exposed to air [42]. It was shown that transfer

Hydrogenated diamond

1.3 eV

Surface acceptors

Ec EVAC

0 eV

EF – 4.2 eV

Ev

Ec UMO +2

UMO +2

UMO +1

UMO +1 LUMO

LUMO EF

e–

EF HOMO OMO-1

FIG. 2.6.6

EVAC

EV

HOMO OMO-1

Transfer doping of H-diamond. The left side refers to neutral diamond and neutral surface acceptors prior to charge exchange. The right side is the band diagram in equilibrium [29]. The zero of the energy scale is the vacuum level. OMO and UMO refer to occupied and unoccupied molecular orbitals, respectively.

210

2. Semiconductor diamond

doping-induced hole conductivity was stable at 300oC when H-terminated diamond was covered by MoO3 and V2O5. More recently, it was reported that when H-terminated diamond was covered by an ReO3 or WO3 layer, the surface conductivity of H-terminated diamond/WO3 showed a hole concentration of 2.18–4.78  1014 cm2 at the very first monolayer coverage (from 1.2 nm). The surface conductivity of the H-terminated diamond/ReO3 was stable up to 450oC with a hole density of 3 1013 cm2 [43]. Interesting MOSFETs are expected by using another oxide such as Al2O3 on the transfer-doped H-terminated diamond covered by an ReO3 layer. When using a high breakdown voltage MOSFET, the double oxide layer should be investigated in the future.

2.6.3 n-Type doping Unlike p-type diamond, n-type diamond is not as easy to achieve. One cannot find a conductive n-type diamond in nature and it must be obtained by incorporating phosphorous atoms into diamond. X-ray absorption fine structure measurement revealed that phosphorus atoms in diamond are located in both substitutional and interstitial sites [44]. Although many efforts have been made in n-type doping by incorporating Li, Na, S, and As atoms into diamond, phosphorous is the only reliable donor up to now [45–47]. Normally, the n-type doping of diamond is accomplished during the epitaxial growth process by adding the PH3 to the gas phase. The first success in n-type doping of diamond was reported in 1997 by Koizumi et al. on the (111) diamond [48]. The Hall mobility was reported to be about 23 cm2/V s at 500 K and the resistivity of the film was 4.43  103 Ω cm at RT. The Hall mobility for a P-doped (111) diamond film with a P concentration of 3  1018 cm3 was improved to 240 cm2/V s at RT [49]. The Hall mobility increased to 660 cm2/V s when the phosphorus concentration was 7  1016 cm3 [50]. The activation energy of phosphorus in diamond deduced from the temperature-dependent conductivity was measured to be 0.57 eV from the conduction band for phosphorus concentration from 2  1015 to 3  1017 cm3. When the phosphorus concentration in the (111)-oriented diamond was reduced, the electron mobility was greatly enhanced. The electron mobility of the film with a phosphorus concentration of 2  1015 cm3 was 1060 cm2/V s at 300 K and 1500 cm2/V s at 225 K [51]. The relationship of mobility and temperature follows the law of T1.5, which is due to acoustic phonon scattering. Most of the efforts in the phosphorous doping of diamond show a relatively low efficiency of 0.1%–3% [49]. By properly designing the substrate holder to change the gas flow, an incorporation efficiency of 10% was achieved from low to high PH3/CH4 flow ratio [52]. One of the problems for the (111)-oriented n-type diamond is the difficulty of mechanical polishing and the limit in size of (111) diamond substrates. Kato et al. reported the n-type doping of homoepitaxial diamond grown on an (001)-oriented diamond substrate by using PH3 as the doping gas [53]. The Hall effect measurements indicate n-type conductivity with an electron mobility of 350 cm2/V s for a phosphorous concentration of 1018 cm3. However, the compensation ratio for (001)-oriented diamond was over 50%, one order of magnitude higher than those of the epilayers on (111)-oriented diamond substrates. The Hall mobility of high-quality (111)-oriented homoepitaxial P-doped diamond films was investigated as a function of doping level for layers with P-density ranging between 6.8 1016 and 6 1019 cm3. The mobility was found to be controlled by the lattice scattering mechanisms in low-doped material below 1017 cm3, by lattice and ionized impurity scattering for a

References

211

moderate doping level from 1017 cm3 to 3 1018 cm3, and by neutral impurity scattering for highly doped material above 1018 cm3. At a high doping level larger than 1019 cm3, hopping conductivity appeared [54]. An activation energy of less than 50 meV was obtained for a P concentration above 1020 cm3, which was due to hopping conductivity. The room temperature resistivity of the heavily phosphorus-doped diamond was 70–150 Ωcm [55, 56].

References [1] C.S. Erasmus, J.P.F. Sellschop, D.M. Bibby, H.W. Fesq, E.J.D. Kable, R.J. Keddy, D.M. Hawkins, D.W. Mingay, S. E. Rasmussen, M.J. Renan, J.I.W. Watterson, Natural diamonds-Major, minor and trace impurities in relation to source and physical propertie, J. Radioanal. Chem. 38 (1977) 133. [2] E. Gaillou, J.E. Post, D. Rost, J.E. Bulter, Boron in natural type IIb blue diamonds: Chemical and spectroscopic measurements, Am. Mineral. 97 (2012) 1. [3] C.M. Breeding, J.E. Shigley, The “type” classfication system of diamond and its importance in gemology, Gems Gemol. 45 (2009) 96–111. ˚ str€ [4] A. Scarani, M. A om, Diamond classification – the diamond types. M&A Gemological Instruments (2015). [5] A.T. Collins, A. Connor, C.H. Ly, A. Shareef, P.M. Spear, High-temperature annealing of optical centers in type-I diamond, J. Appl. Phys. 97 (2005) 083517. [6] S.E. Magan˜aa, T. Ardona, A.M. Zaitsev, LPHT annealing of brown-to-yellow type Ia diamonds, Diam. Relat. Mater. 77 (2017) 159. [7] J. Isberg, G. Ferro, P. Siffert, Diamond Electronic Devices, AIP Conf. Proc. 1292 (2010) 123. [8] E. Kohn, A. Denisenko, Concepts for diamond electronics, Thin Solid Films 515 (2007) 4333. [9] J.P. Lagrange, A. Deneuville, E. Gheeraert, Activation energy in low compensated homoepitaxial boron-doped diamond films, Diam. Relat. Mater. 7 (1998) 1390. [10] D.M. Malta, J.A. von Windheim, H.A. Wynands, B.A. Fox, Comparison of the electrical properties of simultaneously deposited homoepitaxial and polycrystalline diamond films, J. Appl. Phys. 77 (1995) 1536. [11] M.C. Polo, J. Cifre, J. Esteve, Japanese Journal of Applied Physics logo, Vacuum 45 (1994) 1013. [12] S. Yamanaka, H. Watanabe, S. Masai, D. Takeuchi, H. Okushi, K. Kajimura, High-Quality B-Doped Homoepitaxial Diamond Films using Trimethylboron, Jpn. J. Appl. Phys. 37 (1998) L1129. [13] M. Wade, P. Muret, F. Omne`s, A. Deneuville, Technology and electrical properties of ohmic contacts and Schottky diodes on homoepitaxial layers grown on (100) diamond surfaces, Diam. Relat. Mater. 15 (2006) 614. [14] P.N. Volpe, J. Pernot, P. Muret, F. Omne`s, High hole mobility in boron doped diamond for power device applications, Appl. Phys. Lett. 94 (2009) 092102. [15] M.P. Alegre, D. Arau´jo, A. Fiori, J.C. Pinero, F. Lloret, M.P. Villar, P. Achatz, G. Chicot, E. Bustarret, F. Jomard, Critical boron-doping levels for generation of dislocations in synthetic diamond, Appl. Phys. Lett. 105 (2014) 173103. [16] G. Braunstein, R. Kalish, Effective p-type doping of diamond by boron ion implantation, J. Appl. Phys. 54 (1983) 2016. [17] C. Uzan-Saguy, R. Kalish, R. Walker, D.N. Jamieson, S. Prawer, Formation of delta-doped, buried conducting layers in diamond, by high-energy, B-ion implantation, Diam. Relat. Mater. 7 (1998) 1429. [18] T. Vogela, J. Meijer, A. Zaitsev, Highly effective p-type doping of diamond by MeV-ion implantation of boron, Diam. Relat. Mater. 13 (2004) 1822. [19] J.H. Seo, H. Wu, S. Mikael, H. Mi, J.P. Blanchard, G. Venkataramanan, W. Zhou, S. Gong, D. Morgan, Z. Ma, Thermal diffusion boron doping of single- crystal natural diamon, J. Appl. Phys. 119 (2016) 205703. [20] G. Chicot, A. Fiori, P.N. Volpe, T.N. TranThi, J.C. Gerbedoen, J. Bousquet, M.P. Alegre, J.C. Pin˜ero, D. Arau´jo, F. Jomard, A. Soltani, J.C. De Jaeger, J. Morse, J. H€artwig, N. Ranchant, C. Mer-Calfati, J.C. Arnault, J. Delahaye, T. Grenet, D. Eon, F. Omne`s, J. Pernot, E. Bustarret, Electronic and physico-chemical properties of nanometric boron delta-doped diamond structures, J. Appl. Phys. 116 (2014) 083702. [21] T. Kobayashi, T. Ariki, M. Iwabuchi, T. Maki, S. Shikama, S. Suzuki, Analytical studies on multiple delta doping in diamond thin films for efficient hole excitation and conductivity enhancement, J. Appl. Phys. 76 (1994) 1977. [22] J. Scharpf, A. Denisenko, C.I. Pakes, S. Rubanov, A. Bergmaier, G. Dollinger, C. Pietzka, E. Kohn, Transport behaviour of boron delta-doped diamond, Phys. Status Solidi A 210 (2013) 2028. [23] J.E. Butler, A. Vikharev, A. Gorbachev, M. Lobaev, A. Muchnikov, D. Radischev, V. Isaev, V. Chernov, S. Bogdanov, M. Drozdov, E. Demidov, E. Surovegina, V. Shashkin, A. Davidov, H. Tan, L. Meshi, A.

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C H A P T E R

2.7 Diamond semiconductor device and physics

C H A P T E R

2.7.1 Schottky barrier diodes Dan Zhaoa, Juan Wanga, Meiyong Liaob, Hongxing Wanga a

Institute of Wide Band Gap Semiconductors, Xi’an Jiaotong University, Xi’an, P.R. China b Research Center for Functional Materials, National Institute for Materials Science, Tsukuba, Japan

214

2. Semiconductor diamond

2.7.1.1 Surface and metal-diamond interface A barrier is formed at the metal-semiconductor interface when metal contacts the semiconductor. This barrier can control the current conduction and capacitance properties, which are determined by the energetics at the metal-semiconductor interface. The semiconductor electron affinity and the distribution of interface states at the metal-semiconductor junction are key parameters for the barrier height. When no interface states exist at the interface, the barrier height is simply determined by the metal work function for a certain semiconductor. In the other extreme case, the barrier height shows no dependence on the metal work function if there are several interface states. Practically, it is difficult to describe the metal contact to a certain semiconductor by the two simple extreme cases. Sometimes, one can roughly characterize an immediate case between the two extremes by the S parameter to show the linearity variation as the two cases [1]. The electrical properties of the metal-diamond interface are affected by surface terminations, surface morphology, and interface chemistry. The surface structure of diamond depends on the surface terminations and crystallographic orientation. The clean surface diamonds (001) and (111) are not stable and are reconstructed into a 2  1 geometry. The H-terminated diamond (001) surface has a 2  1:2H reconstructed surface. For an O-terminated (001) diamond surface, the surface geometry has a 1 1:O periodicity. For a diamond (111) surface, H-termination leads to the 1  1:H surface structure while for the O-termination, the surface has a 2  1:O reconstruction [2]. In contrast, there are fewer studies on the diamond (110) surface. For diamond (100), the Hterminated diamond surface has a p-type surface conductivity while the oxygen-terminated diamond surface is highly insulating due to its negative polarization. The metal-diamond interface chemistry also plays a key role in determining the interface barrier height [3]. The metals contacting the diamond can be divided into three categories. The first is noncarbide forming metals such as Au, Ag, and Cu. The second is carbide-forming metals such as Al, Ti, Mo, Ta, and V, which react with the diamond at high temperatures [4]. The carbide-forming metals are utilized to fabricate ohmic contact to the diamond. In most cases, Ti is used as the ohmic contact due to the moderate annealing temperature. There is also another group of metals such as Fe, Co, and Ni that have a catalytic effect leading to graphitization of diamond at 450oC. A hydrogen-terminated intrinsic diamond (001) surface exhibits a unique p-type conductivity and a negative electron affinity of about 1.3 eV [5]. The Schottky barrier height depends on the electronegativity or work function of the metal [6]. Mg, Zn, Pb, and Al form a good Schottky contact while Au and Pd are the ohmic contact to the H-terminated diamond [7]. The chemistry at the interface of the O-terminated diamond surface and the metal shows complexity with partially Femi level pining. An oxygen-terminated diamond has a variety of surface states: C-O-C with a positive electron affinity (PEA) of 2.7 eV, hydroxyl(C-OH) with a lower PEA, and ketone (CO), or a combination of all three [8, 9]. The electronic affinity of diamond also depends on the crystal orientation and impurity doping. The electron affinity of the H-terminated heavily P-doped diamond (111) was determined to be 0.2  0.15 eV and reduced to be 0.0  0.15 eV as [P] decreased [10]. The Schottky barrier height of metal to the O-terminated p-type boron-doped diamond (001) is not determined by the metal function or the surface states alone. The values vary from

2.7.1.2 Schottky barrier diode

215

1.2 eV to above 3 eV, depending on the metal species and surface chemistry [11, 12]. A Schottky diode based on an n-type diamond (111) wasn’t investigated as much. The value of the Schottky barrier height is as large as 4.3 eV, showing little dependence on the metals (Al, Ti, Ni, Pt) [13]. To make full use of the high-temperature stability of diamond, a thermally stable metal contact to diamond is certainly necessary. Transition metal carbides such as WC and nitride such as HfN can be utilized [14, 15]. Pin˜ero et al. investigated the chemistry of WC/oxygenterminated diamond and Zr/oxygen-terminated diamond by electron energy loss spectrum (EELS) before and after rannealing [16–18]. It was revealed that thermal treatment at higher temperatures (>600 K) promoted the escape of oxygen for the WC diode while it generated a sharper accumulation of oxygen at the metal/diamond interface for the Zr diode.

2.7.1.2 Schottky barrier diode In the present decade, diamond growth techniques have been improved and doping control methods for p, p +, n-type, and intrinsic diamonds have become available. Therefore, diamond Schottky barrier diodes (SBD) have attracted much research attention to realize high-performance rectifiers providing both high-voltage resistance in reverse operation and low on-resistance in forward operation. A maximum breakdown field of 9.5 MV/cm was reported by analysis of the doping profile and breakdown voltage for a planar Schottky barrier diode (SBD) [19]. Power device capabilities such as a high blocking voltage of Vmax > 10 kV, [20, 21] high current operation at > 20 A, [22] and fast operation with low-loss switching have been recently reported for diamond SBDs. The conventional SBDs with different structures are listed below. A lateral SBD contains ohmic electrodes and Schottky electrodes in the same planar while in the vertical or pseudovertical SBD, a heavily boron-doped diamond substrate or epilayer is used for ohmic contact, as shown in Fig. 2.7.1.1 [23–27]. Ueda et al. reported a high-temperature and high-voltage lateral Cu/diamond SBD [23]. Fig. 2.7.1.2A shows the typical current-voltage (I-V) characteristics of Cu/diamond Schottky diodes measured from room temperature (RT) to 800°C. A high rectification ratio of 105 was observed at RT. The forward current increased with the temperature increasing

FIG. 2.7.1.1 Cross-sectional structure of lateral (left), vertical (middle), and pseudovertical SBDs (right).

216

2. Semiconductor diamond

FIG. 2.7.1.2 (A) I-V characteristics of Cu/diamond Schottky diodes measured from RT to 800°C in vacuum. The inset shows a schematic cross-section and typical top view of the Schottky junctions. (B) Forward I-V characteristics from RT to 600°C. The dotted lines show the results of fitting using the equation for TE mode. The inset shows 1/R versus 1/T [23].

due to the activation of boron acceptors and/or the ohmic resistance value decreased at elevated temperatures. The SBD exhibited rectification characteristics at 600°C with a rectification ratio of 10 [3], although the reverse leakage current increased with the temperature increasing. The SBD can be operated at up to 700°C with a corresponding rectification ratio of 10. The Schottky barrier height (ϕB) and ideality factor n below 600°C were estimated to be 1.4  0.2 eV and 1.7, respectively, by fitting the forward I-V characteristics using the thermionic emission (TE) model (Fig. 2.7.1.2B). The estimated on-resistance R values are less than 1 Ω above 300°C, which were much smaller than the interfacial resistance of the Schottky junctions. The effect of series resistance R can be negligible above 300°C. Fig. 2.7.1.2B plotted reverse resistance 1/R versus temperature 1/T. The activation energy was evaluated to be 0.4 eV, which agrees well with the activation energy of B in diamond semiconductors (0.37 eV). Pseudovertical SBD has also been extensively investigated [28–35]. Traore reported diamond Zr/SBD and ITO/SBD with a pseudovertical architecture. Fig. 2.7.1.3 shows the typical I-V characteristics at RT of diamond Zr/SBD and ITO/SBD. It is clearly seen that Zr/SBD exhibited a good rectification behavior. High current density 103 A/cm2 (at 6 V) and a reverse current density of less than 1  108 A/cm2 up to the maximum voltage (j Vmax j¼1000 V) are obtained. The breakdown (Vbr) of ITO/p-diamond diodes occurred at 200 V and a forward current density of 640 A/cm2 (at 6 V) was obtained. The reverse field (F ¼j Vmax j/d) of 7.7 MV/cm was achieved for Zr/SBD according to the drift layer thickness of 1.3 μm. Vertical SBD provides a facile route for device fabrication [22, 36–44]. A p  layer was grown on a p + layer by the MPCVD system. In this way, the p + layer will act as a contact layer and the p  layer as the drift layer. An ozone treatment was produced by deep UV light to passivate the drift layer surface’s defects. Finally, ohmic and Schottky electrodes were fabricated on the P + and drift layer, respectively. Teraji et al. reported the mechanism of reverse current increase of VSBD. The structure of VSBD is presented schematically in Fig. 2.7.1.4 along with the top view image of the sample. Carbide WC and Ti capped with WC were fabricated on p  and p+ layers, which were used as the Schottky and ohmic electrodes’ metals, respectively. Fig. 2.7.1.5A and B show the

FIG. 2.7.1.3 I-V characteristics of Zr/SBD (samples #1 and #2) and ITO/SBD (sample #1) at RT [28].

WC (Schottky) 100 µm

300 µm

200 µm

P –(NA~ 2 ´ 1015 cm–3)

0.49 µm

P + (NA~1020 cm–3)

500 µm

Ti / WC (ohmic)

FIG. 2.7.1.4 Schematic cross-section of VSBD and top view of the sample [37].

(A)

0

100

150

Voltage [V]

3

4 10–4 10–6 10–8 10–10

Current [A]

10 18 cm – 3 ,f

E

.2x A :7

10–6

b:1

.14

10–4 2 measurement

Current Density [A cm–2]

10–2

10–8

10–12

fb: 0.98 eV

50

Current [A]

10–10

N

TF

10–8

NA:4.5x1014cm–3,

IFL

10–10

10

A :9.8

t ,3 rd en

10–8

E

rem asu

–6 me

10

–6

me

10–4

2

Electrode E2

10–4

asu rem en t x10 16 1 st cm –3 me asu ,f r e : m b 1. ent 18 eV

10–2

2 nd

Current Density [A cm–2]

Electrode E1

1

100

1st measurement

4

N

3

TF

2

1

100

EMAX [MV cm–1]

eV

EMAX [MV cm–1]

NA:4.5x1014cm–3,

IFL

fb: 0.96 eV

10–12

10–10 200

0

(B)

50

100

150

200

Voltage [V]

FIG. 2.7.1.5 Typical reverse J-V characteristics of the (A) stable electrode E1, (B) unstable electrode E2 [37].

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2. Semiconductor diamond

reverse current density-voltage (JR-V) characteristics for a stable electrode E1 and a unstable electrode E2, respectively. The JR-V characteristic changed its slope at the bias voltage of 40– 70 V, which corresponds to the maximum electric field Emax of 0.8–1.4 MV/cm. In the case of silicon SBDs, the image-force lowering (IFL) of the Schottky barrier height is a common mechanism of reverse current increase. The JR including IFL effect is expressed as 0 qffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi1 qφ  q qEmax =4πεD b A, (2.7.1.1) JR ¼ A∗ T2 exp exp @ kT kT where εD is the dielectric constant of diamond. Here, A* of 90 A cm2 K2 and T of RT were used. The barrier height ϕB estimated from the I-V curve was lower than 1.6 eV. At reverse voltage larger than 60 V corresponding to >1.2 MV/cm, the reverse current increased exponentially up to 2  105 A, which was the compliance current of the test system. The reverse bias voltage was applied to the SBDs up to 150 V, which corresponds to Emax of 3.0 MV/cm. The JR-V characteristic of the electrode E1 changed slightly during repetition of the JR-V measurements with the high electric field biasing. The exponential increase part was explained by the thermionic-field emission (TFE) model, which had been observed from SBDs fabricated on other wide-gap semiconductors, such as SiC, GaN, and diamond. Reverse current obeying the TFE model is expressed as  

q 1 (2.7.1.2)  JR ¼ JS exp V kT E0 !1=2 pffiffiffiffiffiffiffiffiffiffiffiffi   A∗ T πqE00 qφb φb   (2.7.1.3) qðV  VP Þ +  exp JS ¼ E0 k cosh 2 qE00 =kT where E0 ¼ E00cosh(E00/kT), E00 ¼ h/4π[NA/m*εD]1/2, VP ¼ EF  EV. EF, EV, and m* represent the Fermi level, the top of the valence band, and the effective mass of electrons, respectively. In addition, E00 reflects the carrier-tunneling probability. The fitted curve is shown by the dashed lines. The NA and ϕB were estimated to be 9.8  1016 cm3 and 1.18 eV, respectively, by using Eqs. (2.7.1.2) and (2.7.1.3). For a different electrode E2, the fitting procedures were the same as those of E1. At a reverse voltage lower than 40 V, NA and ϕB are estimated to be 4.5  1014 cm3 and 0.96 eV, respectively, from fitting by Eq. (2.7.1.1). Next, at a reverse voltage larger than 50 V corresponding to the TFE mode, NA and ϕB are estimated to be 7.2  1016 cm3 and 1.14 eV, respectively, according to Eqs. (2.7.1.2) and (2.7.1.3). A Schottky pn diode (SPND) was also proposed and the principal structure is shown in Fig. 2.7.1.6 [45, 46]. Matsumoto et al. reported a diamond Schottky-pn diode (SPND) using a lightly nitrogen-doped (n ) layer. The schematic and optical images of diamond SPND are shown in Fig. 2.7.1.7A and B, respectively. Ohmic and Schottky electrodes were fabricated on a heavily boron-doped (p +) layer and n  layer, respectively. The impurity concentration and thickness of the n  layer and p + layer measured by secondary ion mass spectrum (SIMS) were 2  1017 cm3 and 80 nm and 2  1020 cm3 and 0.28 mm, respectively. Fig. 2.7.1.8 shows the typical J-V characteristics of the diamond SPND with a Schottky electrode of 50 μm in diameter at RT in air. It is clearly seen that the

FIG. 2.7.1.6

Cross-sectional structure of Schottky pn diode.

FIG. 2.7.1.7

(A) Schematic and (B) optical image of the diamond SPND using the lightly N-doped layer [46].

FIG. 2.7.1.8 (A) Typical J-V characteristics of the diamond SPND. (B) The specific resistance of the diamond SPND estimated using dV/dJ [46].

220

2. Semiconductor diamond

rectification ratio was more than 1013 at 8 V. The forward current density was larger than 20,000 A/cm2 (1 A) at 7 V with a rectification ratio higher than 10 [13]. The breakdown voltage of SPND reached 26 V with a corresponding reverse electric field of 3.3 MV/cm.

References [1] J. Ihm, S.G. Louie, M.L. Cohen, Diamond-Metal Interfaces and the Theory of Schottky Barriers, Phys. Rev. Lett. 40 (1978) 1208. [2] J. Ristein, Surface science of diamond: familiar and amazing, Surf. Sci. 600 (2006) 3677. [3] S. Evans, J.E. Field (Ed.), The Properties of Natural and Synthetic Diamond, Academic Press, London, 1992, pp. 181–214. [4] Y. Koide, M. Yokoba, A. Otsuki, F. Ako, T. Oku, M. Murakami, Diam. Relat. Mater. 6 (1997) 847. [5] B. Rezek, C. Sauerer, C.E. Nebel, M. Stutzmann, J. Ristein, L. Ley, E. Snidero, P. Bergonzo, Fermi level on hydrogen terminated diamond surfaces, Appl. Phys. Lett. 82 (2003) 2266. [6] H. Kawarada, Hydrogen-terminated diamond surfaces and interfaces, Surf. Sci. Rep. 26 (1996) 205. [7] K. Tsugawa, H. Noda, K. Hirose, H. Kawarada, Schottky barrier heights, carrier density, and negative electron affinity of hydrogen-terminated diamond, Phys. Rev. B 81 (2010) 045303. [8] F. Maier, J. Ristein, L. Ley, Electron affinity of plasma-hydrogenated andchemically oxidized diamond (100) surfaces, Phys. Rev. B 64 (2001) 165411. [9] S.J. Sque, R. Jones, P.R. Briddon, Structure, electronics, and interaction ofhydrogen and oxygen on diamond surfaces, Phys. Rev. B 73 (2006) 085313. [10] S. Kono, K. Mizuochi, G. Takyo, N.I. Plusnin, T. Aoyama, T. Goto, T. Abukawa, A. Namba, Y. Nishibayashi, T. Imai, e-J. Surf. Sci. Nanotech. 5 (2007) 33. [11] M. Craciun, C. Saby, P. Muret, A. Deneuville, A 3.4 eVpotential barrier height in Schottky diodes on boron-doped diamond thin films, Diam. Relat. Mater. 13 (2004) 292. [12] M.Y. Liao, Y. Koide, J. Alvarez, T.-s.v.-b.d.p.u.W.C.S. contact, Appl. Phys. Lett. 87 (2005) 022105. [13] M. Suzuki, S. Koizumi, M. Katagiri, T. Ono, N. Sakuma, H. Yoshida, T. Sakai, S. Uchikoga, Electrical characteristics of n-type diamond Schottky diodes and metal/diamond interfaces, Phys. Status Solidi A 203 (2006) 3128. [14] R. Fujii, Y. Gotoh, M.Y. Liao, H. Tsuji, J. Ishikawa, Work function measurement of transition metal nitride and carbide thin films, Vacuum 80 (2006) 832. [15] M.Y. Liao, Y. Koide, J. Alvarez, Thermal stability of diamond photodiodes using WC as Schottky contact, Jpn. J. Appl. Phys. 44 (2005) 7832. [16] P. Muret, A. Traore, A. Marechal, D. Eon, J. Pernot, J.C. Pine˜ro, M.P. Villar, D. Araujo, Potential barrier heights at metal on oxygen-terminated diamond interfaces, J. Appl. Phys. 118 (2015) 204505. [17] J.C. Pin˜ero, D. Araujo, A. Traore, G. Chicot, A. Marechal, P. Muret, M.P. Alegre, M.P. Villar, J. Pernot, Temperature and density dependence metal-oxide-diamond interface investigation by TEM: toward MOS and Schottky power device behavior, Phys. Status Solidi A 211 (2014) 2367–2371. [18] J.C. Pin˜ero, D. Arau´jo, A. Fiori, A. Traore, M.P. Villar, D. Eon, P. Muret, J. Pernot, T. Teraji, Atomic composition of WC/and Zr/O-terminated diamond Schottky interfaces close to ideality, Appl. Surf. Sci. 395 (2017) 200–207. [19] P.N. Volpe, P. Muret, J. Pernot, F. Omnes, T. Teraji, F. Jomard, D. Planson, P. Brosselard, N. Dheilly, B. Vergne, S. Scharnholtz, High breakdown voltage Schottky diodes synthesized on p-type CVD diamond layer, Phys. Status. Solidi A 207 (2010) 2088–2092. [20] P.N. Volpe, P. Muret, J. Pernot, F. Omnes, T. Teraji, Y. Koide, F. Jomard, D. Planson, P. Brosselard, N. Dheilly, B. Vergne, S. Scharnholz, Extreme dielectric strength in boron doped homoepitaxial diamond, Appl. Phys. Lett. 97 (2010) 223501. [21] M. Suzuki, High voltage diamond pin diodes: feasibility study on ultimate properties of diamond toward ultimate power devices, Oyo Buturi 85 (2016) 218–222. [22] V.S. Bormashov, S.A. Terentiev, S.G. Buga, S.A. Tarelkin, A.P. Volkov, D.V. Teteruk, N.V. Kornilov, M. S. Kuznetsov, V.D. Blank, Thin large area vertical Schottky barrier diamond diodes with low on-resistance made by ion-beam assisted lift-off technique, Diam. Relat. Mater. 75 (2017) 78–84.

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[23] K. Ueda, K. Kawamoto, H. Asano, High-temperature and high-voltage characteristics of Cu/diamond Schottky diodes, Diam. Relat. Mater. 57 (2015) 28–31. [24] T. Teraji, M.Y. Liao, Y. Koide, Localized mid-gap-states limited reverse current of diamond Schottky diodes, J. Appl. Phys. 111 (2012) 1104503. [25] A. Fiori, T. Teraji, Y. Koide, Diamond Schottky diodes with ideality factors close to 1, Appl. Phys. Lett. 105 (2014) 133515. [26] Y. Kato, H. Umezawa, S.I. Shikata, X-ray topographic study of defect in p-diamond layer of Schottky barrier diode, Diam. Relat. Mater. 57 (2015) 22–27. [27] H. Umezawa, N. Tatsumi, Y. Kato, S.I. Shikata, Leakage current analysis of diamond Schottky barrier diodes by defect imaging, Diam. Relat. Mater. 40 (2013) 56–59. [28] A. Traore, P. Muret, A. Fiori, D. Eon, E. Gheeraert, J. Pernot, Zr/oxidized diamond interface for high power Schottky diodes, Appl. Phys. Lett. 104 (2014) 052105. [29] H. Umezawa, Y. Mokuno, H. Yamada, A. Chayahara, S. Shikata, Characterization of Schottky barrier diodes on a 0.5-inch single-crystalline CVD diamond wafer, Diam. Relat. Mater. 19 (2010) 208–212. [30] R. Kumaresan, H. Umezawa, S. Shikata, Vertical structure Schottky barrier diode fabrication using insulating diamond substrate, Diam. Relat. Mater. 19 (2010) 1324–1329. [31] R. Kumaresan, H. Umezawa, S. Shikata, Parasitic resistance analysis of pseudovertical structure diamond Schottky barrier diode, Phys. Status Solidi A 207 (2010) 1997–2001. [32] H. Umezawa, S. Shikata, Leakage current analysis of diamond Schottky barrier diodes operated at high temperature, Jpn. J. Appl. Phys. 53 (2014) 04ep04. [33] K. Driche, H. Umezawa, N. Rouger, G. Chicot, E. Gheeraert, Characterization of breakdown behaviour of diamond Schottky barrier diodes using impact ionization coefficients, Jpn. J. Appl. Phys. 56 (2017) 04cr12. [34] D. Zhao, Z.C. Liu, X.F. Zhang, M.H. Zhang, Y.F. Wang, G.Q. Shao, J.W. Zhang, S.W. Fan, W. Wang, H.X. Wang, Analysis of diamond pseudo-vertical Schottky barrier diode through patterning tungsten growth method, Appl. Phys. Lett. 112 (2018) 092102. [35] K. Driche, S. Rugen, N. Kaminski, H. Umezawa, H. Okumura, Electric field distribution using floating guard rings edge-termination for Schottky diodes, Diam. Relat. Mater. 82 (2018) 160–164. [36] H. Umazawa, M. Nagase, Y. Kato, S.I. Shikata, High temperature application of diamond power device, Diam. Relat. Mater. 24 (2012) 201–205. [37] T. Teraji, A. Fiori, N. Kiritani, S. Tanimoto, E. Gheeraert, Y. Koide, Mechanism of reverse current increase of vertical-type diamond Schottky diodes, J. Appl. Phys. 122 (2017) 135304. [38] A. Nawawi, K.J. Tseng, Rusli, G.A.J. Amaratunga, H. Umezawa, S. Shikata, Characterization of vertical Mo/diamond Schottky barrier diode from non-ideal I-V and C-V measurements based on MIS model, Diam. Relat. Mater. 35 (2013) 1–6. [39] A. Nawawi, K.J. Tseng, Rusli, G.A.J. Amaratunga, H. Umezawa, S. Shikata, Design and optimization of planar mesa termination for diamond Schottky barrier diodes, Diam. Relat. Mater. 36 (2013) 51–57. [40] H. Umezawa, Y. Kato, S.I. Shikata, 1Ω on-resistance diamond vertical-Schottky barrier diode operated at 250 °C, Appl. Phys. Express 6 (2013) 011302. [41] M. Nagase, H. Umezawa, S. Shikata, Vertical diamond Schottky barrier diode fabricated on insulating diamond substrate using deep etching technique, IEEE Trans. Electron Devices 60 (2013) 1416–1420. [42] H. Umezawa, S. Shikata, T. Funaki, Diamond Schottky barrier diode for high-temperature, high-power, and fast switching applications, Jpn. J. Appl. Phys. 53 (2014) 05fp06. [43] V.D. Blank, V.S. Bormashov, S.A. Tarelkin, S.G. Buga, M.S. Kuznetsov, D.V. Teteruk, N.V. Kornilov, S. A. Terentiev, A.P. Volkov, Diam. Relat. Mater. 57 (2015) 32–36. [44] D. Zhao, C. Hu, Z.C. Liu, H.X. Wang, W. Wang, J.W. Zhang, Diamond, MIP structure Schottky diode with different drift layer thickness, Diam. Relat. Mater. 73 (2017) 15–18. [45] T. Makino, H. Kato, N. Tokuda, M. Ogura, D. Takeuchi, K. Oyama, S. Tanimoto, H. Okushi, S. Yamasaki, Diamond Schottky-pn diode without trade-off relationship between on-resistance and blocking voltage, Phys. Status Solidi A 207 (2010) 2105–2109. [46] T. Matsumoto, T. Mukose, T. Makino, D. Takeuchi, S. Yamasaki, T. Inokuma, N. Tokuda, Diamond Schottky-pn diode using lightly nitrogen-doped layer, Diam. Relat. Mater. 75 (2017) 152–154.

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C H A P T E R

2.7.2 Metal-Semiconductor Field-Effect Transistors Dan Zhaoa, Juan Wanga, Masataka Imurab, Hongxing Wanga a

Institute of Wide Band Gap Semiconductors, Xi’an Jiaotong University, Xi’an, P.R. China b Research Center for Functional Materials, National Institute for Materials Science (NIMS), Japan

Wide bandgap semiconductor materials provide the solution for the requirements of wireless communication systems: high RF output power density and high operation voltages with low energy consumption [1]. As the widest bandgap semiconductor (5.47 eV), diamond exhibits superior electrical and thermal properties such as high breakdown field strength (>10 MV/cm), high intrinsic carrier mobility (4500cm2/V s for electrons and 3800 cm2/V s for holes), high saturation velocity (1.5  107cm/s for electrons and 1.1  107cm/s for holes), low dielectric constant (5.7), and very high thermal conductivity (22W/cm K) [2–4], making it a promising candidate for the application of high-frequency, high-power, and hightemperature electric devices. Intrinsic diamond is an insulating material, and the conductivity is extremely low. Nevertheless, diamond has unique surface properties. Hydrogen-terminated diamond shows a ptype conductive channel with a sheet charge concentration of about 1 1013 cm2 in the air. So far, several models have been proposed to explain the formation of the hole channel such as H-induced acceptor-type surface states, charged gas molecules adsorbed on diamond surface, electrochemical model involving oxonium ions, and a charge transfer model employing an aqueous surface layer [5, 6], which are all controversial. Hydrogen-terminated diamond can be used to fabricate field-effective transistors (FET). Kasu et al. achieved a maximum RF output power density of 2.1 W/mm at 1 GHz [7] and cut-off frequencies fT and fMAX of 45 and 120 GHz [8].

2.7.2.1 NO2 adsorption to increase sheet hole density The surface conductivity of the H-terminated diamond surface decreases in vacuum but returns to the initial value reexposed to air, which indicates that the adsorbents in the air are essential to form the hole channel. It is demonstrated that the highest conductivity can be obtained exposed to the NO2 gas among the many gases in the air [5]. The sheet hole concentration (ps) significantly increased up to 1.3 1014 cm2 at 300 ppm NO2, one order of

2.7.2.2 DC characteristics of H-terminated diamond FETs

223

magnitude higher than previously reported values. It was believed that the increase of the hole sheet density was due to the decomposition of NO2 molecules into N2 and O2. The resultant maximum drain current and the transconductance markedly increased by 1.8 times and 1.5 times, respectively, after exposure to NO2. It is noticed that the NO2 adsorbates was unstable, which leads to the instability problem of hydrogen-terminated FET characteristics. The thermal instability of the H-terminated diamond surface hinders the practical application of diamond power transistors. Accordingly, it is necessary to find an effective method to passivate NO2 adsorbates. Kasu et al. report an atomic-layer-deposited (ALD) Al2O3 overlayer deposition method, which made the NO2-adsorbed H-terminated diamond surface more stable and kept the sheet hole concentration and mobility high, even at higher temperatures [9]. This thermal stabilization technique realized thermally stable H-terminated diamond FETs with a reproducible high drain current and almost no degradation. It was revealed with with Al2O3 passivation, ps kept nearly invariable at 423 K (150°C) for 24 h in vacuum. On the contrast, the ps of the sample without the Al2O3 overlayer sharply decreased by an order of magnitude in 10 min, which indicated the thermal desorption of NO2 molecules at 150°C.

2.7.2.2 DC characteristics of H-terminated diamond FETs Using the NO2 adsorption and Al2O3 passivation technique, Hirama et al. improved the drain current (IDS) of hydrogen-terminated diamond field effective-field transistors (FETs) and obtained the maximum drain current density of 1.35 A/mm [10]. Fig. 2.7.2.1 shows the drain current-voltage (IDS-VDS) characteristics of the H-terminated diamond MESFET with a 0.4-μm gate length by using the NO2 adsorption and Al2O3 passivation technique. Gate voltage (VGS) varied from +11 to 5 V in steps of 2 V. The threshold voltage was +11 V, and FIG. 2.7.2.1 IDS-VDS characteristics of a passivated diamond FET with LG ¼0.4 μm (VGS ¼+11 to 5 V; ΔVGS ¼ 2 V) [10].

224

2. Semiconductor diamond

no drain leakage current was observed. When VGS was 5 V, the drain current density reached the maximum of 1.35 A/mm. This value was the highest IDSmax ever reported for H-terminated diamond FETs. In the saturation region, a decrease of IDS was not observed, which indicates that the influence of the self-heating effect on the drain current was very small, owing to the high thermal conductivity of diamond.

2.7.2.3 RF characteristics of H-terminated diamond FETs Fig. 2.7.2.2 shows the frequency dependence of the current gain, j h21 j2, and the unilateral power gain, U. From the gain values at lower frequency, the cut-off frequencies of fT and fMAX were extrapolated to be 35 and 70 GHz, respectively [11–13]. The maximum RF output power density, maximum power gain, and power-added efficiency (PAE) at a VGS of 1.0 V and VDS of 25 V were 2 W/mm, 18 dB, and 33%, respectively.

2.7.2.4 Thermally stable operation of H-terminated diamond FETs Hirama et al. applied the NO2 adsorption and Al2O3 passivation technique to fabricate Hterminated diamond FETs. It was demonstrated that the IDS kept constant for more than 1 h with no increase of gate leakage current at 200°C in vacuum (3 103 Pa). Fig. 2.7.2.3 shows the IDS-VDS characteristics of a passivated diamond FET with a 0.4-μm gate length at 200°C and room temperature before and after the 200°C heating cycle. VGS varied from +2 V to 4 V in steps of 1 V. At room temperature before heating, the maximum drain current density IDSmax was 200 mA/mm and then decreased to 180 mA/mm at 200°C. After the 200°C heating cycle, IDSmax returned to the original value of 200 mA/mm. It was observed that IDS at every VGS agreed well with each other at room temperature before and after 200°C FIG. 2.7.2.2 RF small-signal characteristics of a passivated diamond FET. The cut-off frequencies of fT (35 GHz) and fMAX (70 GHz) were extrapolated [11].

40 LG : 0.1 mm WG : 100 mm

H21 2 MSG

Gain (dB)

30 fT : 35 GHz fmax : 70 GHz 20

10

fmax fT

0 1

10 Frequency (GHz)

100

References

225

FIG. 2.7.2.3 IDS-VDS characteristics of a 0.4-μm gate-length diamond FET with a passivation layer before, during, and after 200°C heating [11].

heating. The thermal stability of H-terminated diamond FETs was improved by the NO2 adsorption and Al2O3 passivation technique. The technique may be able to open a new path to high-temperature applications of diamond FETs.

2.7.2.5 Conclusions A diamond semiconductor is a promising candidate for high-power, high-frequency, and high-temperature applications due to its remarkable properties. The current status of diamond diodes has been introduced. The performance of diamond diodes has been improved since the establishment of homoepitaxial growth techniques and doping control. New processing techniques to form edge-termination structures, especially on selectively doped substrates using ion implantation or selective area growth, are required for effective use of the attractive properties of diamond. A deeper understanding of surface, interface, and defect structures is necessary in order to improve both device fabrication and performance. For H-terminated diamond FETs with a hole conductive channel, there remains the thermal stability problem. To improve the hole density and thermal stability, NO2 adsorption and the Al2O3 passivation technique were adopted. Stable FET operation at 200°C in a vacuum was demonstrated.

References [1] P. Calvani, G. Conte, D. Dominijanni, et al., Hydrogen terminated diamond MESFETs: new technology for RF power applications, in: The 5th European Microwave Integrated Circuits Conference, IEEE, 2010, pp. 122–125. [2] J. Isberg, J. Hammersberg, E. Johansson, et al., High carrier mobility in single-crystal plasma-deposited diamond, Science 297 (2002) 1670–1672. [3] L.S. Pan, D.R. Kania, Diamond: Electronic Properties and Applications, Springer, United States, 1995. [4] L. Reggiani, S. Bosi, C. Canali, et al., Hole-drift velocity in natural diamond, Phys. Rev. B 23 (1981) 3050–3057.

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[5] M. Kubovic, Improvement of hydrogen-terminated diamond field effect transistors in nitrogen dioxide atmosphere, Appl. Phys. Express 2 (2009) 6502. [6] D. Petrini, K. Larsson, Structural and energetic consideration of H, F, S, Cl-terminated Diamond (100) and (111) surfaces: quantum mechanical study, J. Phys. Chem. C (2007). [7] M. Kasu, K. Ueda, H. Ye, et al., 2 W/mm output power density at 1 GHz for diamond FETs, Electron. Lett. 41 (2005) 1249–1250. [8] K. Ueda, M. Kasu, Y. Yamauchi, et al., Diamond FET using high-quality polycrystalline diamond with f/sub T/ of 45 GHz and f/sub max/ of 120 GHz, IEEE Electron Device Lett. 27 (2006) 570–572. [9] M. Kasu, H. Sato, K. Hirama, Thermal stabilization of hole channel on H-terminated diamond surface by using atomic-layer-deposited Al2O3 overlayer and its electric properties, Appl. Phys. Express 5 (2012) 5701. [10] K. Hirama, H. Sato, Y. Harada, et al., Diamond field-effect transistors with 1.3A/mm drain current density by Al2O3 passivation layer, Jpn. J. Appl. Phys. 51 (2012) 080112. [11] K. Hirama, H. Sato, Y. Harada, et al., Thermally stable operation of H-terminated diamond FETs by, adsorption and, passivation, IEEE Electron Device Lett. 33 (2012) 1111–1113. [12] M. Kasu, T. Oishi, Diamond RF power transistors: present status and challenges, in: European Microwave Integrated Circuit Conference, IEEE, 2014, pp. 146–148. [13] M. Kasu, T. Oishi, Recent progress of diamond devices for RF applications, in: Compound Semiconductor Integrated Circuit Symposium, IEEE, 2016, pp. 1–4.

C H A P T E R

2.7.3 Metal-oxide/diamond interface and MOSFETs Yanfeng Wang, Zhangcheng Liu, Hongxing Wang Institute of Wide Band Gap Semiconductors, Xi’an Jiaotong University, Xi’an, P.R. China

2.7.3.1 MOSFETs based on H-terminated diamond surface To realize the true diamond MOSFET properties, both a sufficient carrier density and a reliable MOS interface compatible with extreme diamond properties are required. Most of the efforts on diamond MOSFETs have been focused on the H-terminated surface [1–10]. Fig. 2.7.3.1 shows the cross-sectional structure of the H-diamond MOSFET fabricated by Kawarada et al. [1] The detailed fabrication process was shown in Fig. 2.7.3.2. A 500 nm thickness H-diamond was deposited on a type Ib (001) diamond substrate, as shown in Fig. 2.7.3.2A. The diamond surfaces were oxidized by UV ozone treatment to form a partially C–O bonded surface, as shown in Fig. 2.7.3.2B, which was sufficiently insulating for the isolation of the conducting channel composed of a C–H bonded surface [11]. This is because there was a

2.7.3.1 MOSFETs based on H-terminated diamond surface

227

FIG. 2.7.3.1 The cross-sectional structure of the H-diamond MOSFET [1].

FIG. 2.7.3.2 Fabrication process of the H-diamond MOSFET [1].

potential barrier of more than 2.0 eV between the C–H and C–O surfaces due to the difference in electron affinity [12, 13]. Au/Ti contacts for the source and drain were formed by the sequential deposition of the two metals on the diamond surface. The Au/Ti on the partially C–O bonded surface was annealed for 1 h in H2 at 600°C to form a thin TiC layer, as shown in Fig. 2.7.3.2C [14]. Areas other than the channel and parts of the source and drain were masked by photoresistance. H2 plasma exposure was then performed through the mask pattern of photoresistance to form a C–H bonded surface (channel) bounded laterally by the source contact, drain contact, and C–O surface for isolation, as shown in Fig. 2.7.3.2D. After the photoresistance was removed, Al2O3 was deposited on the whole surface at 450°C by alternately supplying trimethyl aluminum and H2O [15]. This Al2O3was used as a gate insulator and the passivation layer for a C–H bonded surface channel. The thickness of the Al2O3 layers was controlled from 10 to 200 nm. On the source and drain contact regions, the deposited Al2O3 was partially removed to form an electrical contact, as shown in Fig.2.7.3.2E. Finally, a metal gate electrode (Al) was formed on the Al2O3, as shown in Fig. 2.7.3.2F. To fabricate MOSFET based on an H-terminated diamond surface, two common points are noted as follows. First, to enhance the performance of H-diamond MOSFETs, an excellent ohmic contact between the source/drain and the H-diamond is needed [16–20]. Until now, many investigations of ohmic contact between metal and diamond have been reported, including metals such as Au, Pt, Ir, Ti, and palladium (Pd) [9, 10, 21–23]. However, Au could peel off during the fabrication process, which indicates quite poor adherence of Au on the H-

228

2. Semiconductor diamond

diamond film. Pt and Pd could be corroded in harsh environments that cannot adapt to the application of diamond electronic devices [24]. The second is the isolation of H-terminated diamond MOSFETs. Thanks to the insulation of C–O diamond surfaces and the high potential barrier of C–H and C–O diamond surfaces, C–O surfaces can be used to isolate the C–H conducting channel. Many methods can be used to form a C–O surface, such as UV/ozone, reactive ion etching, etc.

2.7.3.2 Normally-on H-diamond MOSFET After diamond is treated by hydrogen plasma, 2DHG will be formed on the diamond surface. The 2DHG can be directly used as a MOSFET channel, indicating that normally on Hdiamond MOSFETs are easy to be fabricated. A typical work on normally on H-diamond MOSFETs is from Kawarada [1]. Fig. 2.7.3.3 shows the IDS-VDS characteristics of H-diamond MOSFETs with an Al2O3 dielectric layer at (a) room temperature and (b) 400°C. The 10 nm Al2O3 dielectric layer was deposited by atomic layer deposition at 450°C. In Fig. 2.7.3.3, a typical IDS-VDS characteristic showed saturation curves due to pinch-off with a clear cut-off region. At 400°C, IDS-VDS characteristics did not differ substantially from those obtained at room temperature. The saturation behavior was almost perfect. The maximum transconductance normalized by the gate width for the MOSFETs with an Al2O3 thickness of 10 nm was 9–10 mS/mm. The maximum drain current density normalized by the gate width was 30–45 mA/mm, which was saturated at high gate voltages. Fig. 2.7.3.4 shows the breakdown characteristics of H-diamond MOSFETs with a 200 nm Al2O3 dielectric layer. A high breakdown voltage was obtained in FET with the 200 nm-thick Al2O3 gate oxide without a field plate. The maximum breakdown voltage VB was 606 V at a gate-to-drain distance LGD of 7 μm. The averaged electric field VB/LGD was 120 V/μm in this case. VB/LGD was often used as a measure (criterion) for high-voltage durability in planar FETs, with 100 V/μm being a critical value for lateral power devices. However, the

FIG. 2.7.3.3

IDS-VDS characteristics of H-diamond MOSFETs with an Al2O3 dielectric layer at (A) room temperature and (B) 400°C [1].

2.7.3.4 MOSFET based on boron-doped diamond

229

FIG. 2.7.3.4 Breakdown characteristics of H-diamond MOSFETs [1].

breakdown voltage of more than 1000 V obtained in lateral SiC MOSFETs VB/LGD was less than 100 V/μm for these FETs, and slightly less than the value of AlGaN/GaN FET [47, 48]. Recently, an AlGaN/AlGaN FET with a lateral breakdown field of 160 V/μm was reported, which was one of the highest values recorded for planar FETs [25]. By properly designing the value of LGD, a high breakdown voltage of 1700 V was achieved [7].

2.7.3.3 Normally-off H-diamond MOSFET Fail-safe systems strongly require normally off operation in H-diamond MOSFETs. Recently, several groups reported that normally off MOSFETs have been fabricated in the diamond field. Kitabayashi et al. successfully realized a normally off H-diamond MOSFET in which the Hterminated conductive channel under the gate electrode was partially treated by UV/ozone to form partially oxygen-terminated nonconductive areas [26]. The dielectric layer of an Al2O3 film with a thickness of 200 nm was deposited by ALD at 450°C. Umezawa et al. also used a partially oxygen-terminated channel to realize a MOSFET with a normally off property [27]. A double-layer oxide such as an LaAlO3/Al2O3 dielectric also led to the normally off Hdiamond MOSFETs by annealing the sample at 180°C for 10 min, which is likely due to the imperfect nature of the insulators [28]. Wang et al. successfully fabricated normally off Hdiamond MOSFETs with a 3-nm Al2O3 dielectric layer formed by thermal oxidation of Al [29].

2.7.3.4 MOSFET based on boron-doped diamond Boron-doped diamond was also applied for p-type diamond MOSFET fabrication. Recently, Pham et al. reported a lateral diamond MOSFET based on a boron-doped

230

2. Semiconductor diamond

GAT

LGS Ti/Pt/Au Tic

LG Ti/Pt/Au SCR

p-type diamond h

MOSFET fingers

DRAIN

E

RCE SOU

diamond with a high breakdown voltage and breakdown electric field [30]. Due to the deep level of boron, the deep depletion concept was utilized for the operation of the boron-doped diamond MOSFET. Fig. 2.7.3.5 shows the conceptual device structure and optical image of the MOSFETs. The ohmic contacts were deposited on a hydrogen termination surface and followed by an annealing process at 500°C. Before oxide deposition, the surface hydrogen termination was changed into oxygen termination through deep UV ozone treatment. The thickness of the p-type monocrystalline channel was 190 nm. The boron concentration was 1.75  1017 cm3 and the thickness of the ALD-Al2O3 was 20 nm deposited at 380oC. The typical current-voltage characteristics of the p-type diamond MOSFET without any oxide annealing at high temperature are displayed in Fig. 2.7.3.6A. The transistor is

LGD

w

100mm Ti/Pt/Au AI2O3 Tic

I

Space Charge Region (SCR)

(A) Substrate HPHT [100] lB

MOS Capacitors

d Corbino MOSFETs

(B)

FIG. 2.7.3.5 (A) Conceptual cross-section structure of the depletion mode oxygen-terminated boron-doped diamond MOSFET. (B) Optical image of the top view of the boron-doped diamond MOSFET [30].

FIG. 2.7.3.6 properties [30].

I-V characteristics of the p-type diamond MOSFET (A) without and (B) with transfer electrical

References

231

normally-on with a threshold voltage in the deep depletion regime of +7 V, a maximum saturation current of 1.9 μA/mm at VSD ¼ 10 V and VGS ¼5 V, at room temperature. A high hole mobility (1000  200cm2/(V s)) was evaluated. The existence of oxygen termination introduced the Fermi-Level-Pinning effect and no accumulation regime appeared under the gate. When an oxide annealing process was applied, the accumulation regime was observed for negative gate bias. Before the experimental voltage breakdown at 200 V, the gate leakage was below 0.6 nA/ mm at room temperature and the peak electric field in diamond at the gate edge was simulated to be at 4 MV/cm. Due to the significantly high activation energy of boron-doped diamond, the performances of such MOSFETs will be improved at higher temperatures. The IDS-VDS characteristics from RT to 250°C shows an improvement of 25 times for the saturation currents. Better boron dopant ionization has a stronger effect at high temperatures than the decrease of carrier mobility. Overall, the specific ON state resistance of such diamond p-type MOSFETs can be very attractive at high junction temperatures.

2.7.3.5 Conclusion Diamond p-channel MOSFETs were demonstrated on both an H-terminated diamond surface and a boron-doped diamond. High temperature operation up to 400oC and high breakdown voltage up to 1700 kV were achieved on an H-terminated diamond surface channel MOSFET. The maximum drain current can reach more than 100 mA/mm. For the gate insulators, ALD-Al2O3 shows the best performance and reliability among the reported oxides. For the boron-doped diamond MOSFET, deep depletion mode was utilized and the cleat transistor behavior was observed. Despite the progress, the interface stability of the H-terminated diamond MOS should be investigated, and there is still a far distance to go to achieve a high-current boron-doped transistor.

References [1] H. Kawarada, H. Tsuboi, T. Naruo, T. Yamada, D. Xu, A. Daicho, T. Saito, A. Hiraiwa, C-H surface diamond field effect transistors for high temperature (400 °C) and high voltage (500 V) operation, Appl. Phys. Lett. 105 (2014) 013510. [2] H. Kawarada, High-current metal oxide semiconductor field-effect transistors on H-terminated diamond surfaces and their high-frequency operation, Jpn. J. Appl. Phys. 51 (2012) 090111. [3] T. Saito, K.H. Park, K. Hirama, H. Umezawa, M. Satoh, H. Kawarada, Z.Q. Liu, K. Mitsuishi, K. Furuya, H. Okushi, Fabrication of metal–oxide–diamond field-effect transistors with submicron-sized gate length on boron-doped (111) H-terminated surfaces using electron beam evaporated SiO2 and Al2O3, J. Electron. Mater. 40 (2011) 247–252. [4] K. Hirama, S. Miyamoto, H. Matsudaira, K. Yamada, H. Kawarada, T. Chikyo, H. Koinuma, K. Hasegawa, H. Umezawa, Characterization of diamond metal-insulator-semiconductor field-effect transistors with aluminum oxide gate insulator, Appl. Phys. Lett. 88 (2006) 112117. [5] J. Zhao, J.W. Liu, L.W. Sang, M.Y. Liao, D. Coathup, M. Imura, B.G. Shi, C.Z. Gu, Y. Koide, H. Ye, Assembly of a high-dielectric constant thin TiOx layer directly on H-terminated semiconductor diamond, Appl. Phys. Lett. 108 (2016) 012105. [6] K. Hirama, H. Takayanagi, S. Yamauchi, High-performance p-channel diamond MOSFETs with alumina gate insulator, in: Electron Devices Meeting, 2007 (IEDM 2007), IEEE International, 2008, pp. 873–876. [7] H. Kawarada, T. Yamada, D. Xu, Y. Kitabayashi, M. Shibata, D. Matsumura, M. Kobayashi, T. Saito, T. Kudo, M. Inaba, A. Hiraiwa, Diamond MOSFETs using 2D Hole Gas with 1700V Breakdown Voltage,

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[8] [9] [10] [11] [12] [13]

[14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28]

[29]

[30]

2. Semiconductor diamond

in: Proceedings of the 2016 28th International Symposium on Power Semiconductor Devices and ICs (ISPSD), June 12–16, Prague, Czech Republic, 2016. W. Wang, K. Fu, C. Hu, F.N. Li, Z.C. Liu, S.Y. Li, F. Lin, J. Fu, J.J. Wang, H.X. Wang, Diamond based field-effect transistors with SiNx and ZrO2 double dielectric layers, Diam. Relat. Mater. 69 (2016) 237–240. W. Wang, C. Hu, F.N. Li, S.Y. Li, Z.C. Liu, F. Wang, J. Fu, H.H. Wang, Palladium ohmic contact on hydrogenterminated single-crystal diamond film, Diam. Relat. Mater. 63 (2016) 175–179. J.W. Liu, Y. Koide, Fabrication of hydrogenated diamond metal-insulator-semiconductor field-effect transistors, Methods Mol. Biol. 1572 (2017) 217–232. T. Sakai, K.S. Song, H. Kanazawa, Y. Nakamura, H. Umezawa, M. Tachiki, H. Kawarada, Ozone-treated channel diamond field-effect transistors, Diam. Relat. Mater. 12 (2003) 1971. J.B. Cui, J. Ristein, L. Ley, Electron affinity of the bare and hydrogen covered single-crystal diamond (111) surface, Phys. Rev. Lett. 81 (1998) 429. M. Tachiki, Y. Kaibara, Y. Sumikawa, M. Shigeno, H. Kanazawa, T. Banno, K.S. Song, H. Umezawa, H. Kawarada, Characterization of locally modified diamond surface using Kelvin probe force microscope, Surf. Sci. 581 (2005) 207. Y. Jingu, K. Hirama, H. Kawarada, Ultrashallow TiC source/drain contacts in diamond MOSFETs formed by hydrogenation-last approach, IEEE Trans. Electron Devices 57 (2010) 966. A. Hiraiwa, A. Daicho, S. Kurihara, Y. Yokoyama, H. Kawarada, Refractory two-dimensional hole gas on hydrogenated diamond surface, J. Appl. Phys. 112 (2012) 124504. C.M. Zhen, X.Q. Wang, X.C. Wu, C.X. Liu, D.L. Hou, Au/p-diamond ohmic contacts deposited by RF sputtering, Appl. Surf. Sci. 255 (2008) 2916–2919. K.L. Moazed, R. Nguyen, J.R. Zeidler, Ohmic Contacts to semiconducting diamond, IEEE Electron Device Lett. 9 (1988) 350–354. M. Werner, How to fabricate low-resistance metal-diamond contacts, Diam. Relat. Mater. 5 (1996) 723–727. M. Werner, Very low resistivity A1-Si ohmic contacts to boron-doped polycrystalline diamond films, Diam. Relat. Mater. 3 (1994) 983–985. M. Yokoba, Carrier transport mechanism of Ohmic contact to p-type diamond, J. Appl. Phys. 81 (1997) 6815. C. Verona, W. Ciccognani, S. Colangeli, F.D. Pietrantonio, Gate–source distance scaling effects in H-terminated diamond MESFETs, IEEE Trans. Electron Devices 62 (2015) 1150. K. Tsugawa, H. Noda, K. Hirose, H. Kawarada, Schottky barrier heights, carrier density, and negative electron affinity of hydrogen-terminated diamond, Phys. Rev. B 81 (2010) 045303. Y.F. Wang, Ohmic contact between iridium film and hydrogen-terminated single-crystal diamond, Sci. Rep. 7 (2017) 12157. K. Fumihiro, Electrochemical corrosion of platinum electrode in concentrated sulfuric acid, J. Power Sources 172 (2007) 698–703. T. Nanjo, A. Imai, Y. Suzuki, Y. Abe, T. Oishi, M. Suita, E. Yagyu, Y. Tokuda, AlGaN channel HEMT with extremely high breakdown voltage, IEEE Trans. Electron Devices 60 (2013) 1046. Y. Kitabayashi, Normally-off C-H diamond MOSFETs with partial C-O Channel achieving 2-kV breakdown voltage, IEEE Electron Device Lett. 38 (3) (2017). H. Umezawa, RF diamond transistors: current status and future prospects, Jpn. J. Appl. Phys. 44 (2005) 7789–7794. J.W. Liu, M.Y. Liao, M. Imura, T. Matsumoto, N. Shibata, Y. Ikuhara, Y. Koide, Control of normally on/off characteristics in hydrogenated diamond metal-insulator-semiconductor field-effect transistors, J. Appl. Phys. 118 (2015) 115704. Y.F. Wang, X.H. Chang, X.F. Zhang, J. Fu, S.W. Fan, R.A. Bu, J.W. Zhang, W. Wang, H.X. Wang, J.J. Wang, Normally-off hydrogen-terminated diamond field-effect transistor with Al2O3 dielectric layer formed by thermal oxidation of Al, Diam. Relat. Mater. 81 (2018) 113–117. T.T. Pham, J. Pernot, G. Perez, D. Eon, E. Gheeraert, N. Rouger, Deep-depletion mode boron-doped monocrystalline diamond metal oxide semiconductor field effect transistor, IEEE Electron Device Lett. 38 (2017) 1571.

Further reading [31] T.T. Pham, N. Rouger, C. Masante, G. Chicot, F. Udrea, D. Eon, E. Gheeraert, J. Pernot, Deep depletion concept for diamond MOSFET, Appl. Phys. Lett. 111 (2017) 173503.

2.7.4 Junction FETs

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2.7.4 Junction FETs Zhangcheng Liu, Hongxing Wang Institute of Wide Band Gap Semiconductors, Xi’an Jiaotong University, Xi’an, P.R. China

Up to now, no shallow n-type doping has been available in diamond. As a result, all diamond-based FETs are p-channel devices using either boron doping or an H-terminated surface conductive layer as the channel. Even with electrical activation of the boron dopant at ambient temperature, researchers have found that with increasing doping concentration, the activation energy of boron decreases [1, 2] so that full activation at room temperature can be obtained for an effective peak doping in the range of 1020 cm3. This opens up a possibility to fabricate a device with full charge activation at room temperature. However, in order to realize a full charge modulation by the gate diode within the materials breakdown limit of 3 106–1 107 V/cm, the total channel sheet charge should be kept below the order of 1013 cm2 [3, 4]. Therefore the width of the doped region has to be in the order of 1–2 nm. This leads to boron δ-doping profiles. Obtaining such profiles is a challenge because there is still significant residual doping present in the layer above the pulse. Elastic recoil detection measurements and capacitance-voltage (CV) profiling indicated parasitic doping tails toward the surface, which may prevent the modulation of the channel charge by the gate and may cause degradation of the Schottky contact characteristics. To solve this problem, nitrogen doping was applied, leading to the formation of a nitrogen-boron pn-junction because nitrogen in diamond was confirmed to be a deep donor [5]. Therefore, it seems attractive and possible to fabricate devices with δ-doped channels where the control diode contains such a pn-junction. The potential advantages of diamond pn-junctions in diamond δ-FET structures over Schottky contacts are the low leakage current and high breakdown as well as the higher built-in voltage of above 3.0 V compared to the 1.7 V of the Schottky contact [6]. This leads to a shift in the pinch-off voltage toward enhancement-mode operation. In the junction FET (JFET), the gate metal is separated from the channel by an N-doped layer, as shown in Fig. 2.7.4.1. The n-doped layer is a lossy dielectric represented by a leakage resistor in parallel to the dielectric capacitance inserted in between the gate metal and channel. Depending on the layer properties and dimensions, the behavior can be either resistive or capacitive. Depending on the position of the gate metal, two types of JFETs—back-gate JFETs (BGFETs) and top-gate JFETs (TGFETs)—were proposed, as shown in Fig. 2.7.4.2 [7]. In BGFETs, the nitrogen-doped Ib substrate was used as a large area gate underneath the channel. The substrate material has to be highly n-doped. Additionally, the n+-doped substrate also serves as the gate contact. These devices are not expected to operate at room temperature

234

2. Semiconductor diamond

FIG. 2.7.4.1 N-doped lossy dielectric layer between gate metal and channel of a diamond JFET [6].

Gate metal

RD

N-doped lossy dielectric

CD

dD

Depletion region

CPN Channel

Schematic cross-section of δ-doped channel FETs: (A) structure with the gate at the back (BGFET), (B) structure with the gate at the top (TGFET) [7].

FIG. 2.7.4.2

Source

Drain

d-doped layer Transition layer (nominally undoped)

lb (N-doped) Diamond substrate & gate

(A)

gate p+

- doped contact layer

Source

Gate

Drain

d-doped layer

(B)

Buffer layer (nominally undoped)

N-doped layer (n+ - and transition layer)

lb (N-doped) Diamond substrate

due to the fact that the nitrogen donor with an ED ¼ 1.7 eV [1] is frozen out, and therefore a sufficient conductivity of the substrate is expected only at elevated temperatures. In this case, the substrate acts as a series resistance to the pn-junction. In TGFETs, the n+-control layer is grown on top of the δ-layer. For the case of a narrow control layer, the heavily n+-doped part of the resulting junction represents a semiinsulating, lossy dielectric at low temperature and high frequency. In this case, the gate contact is represented by the gate metal contact, resulting in a possibility of operation at room temperature. The control layers in this structure also serve as separation layers for gate metallization. The DC output characteristics of the two types of JFETs are shown in Fig. 2.7.4.3. The gate width of the BGFET was 250 μm and the gate length was 2 mm. The effective gate length for devices with the gate at the back equals the source to drain separation, that is, the length of the channel. As expected, at room temperature only weak modulation of the drain current was

235

2.7.4 Junction FETs

–20

–120

VG= –10V –15

DVG= 5V –40

ID (mA/mm)

ID (mA/mm)

IG = 2µm WG = 250µm dsl = 320nm –80 T = 250°C

VG= –4V IG = 20µm WG = 100µm T = 200°C DVG= 2V

–10

–5 VG= 30V 0

0

(A)

–10 VD(V)

0 0

–20

(B)

VG= 22V –10

–20

–30

VD(V)

FIG. 2.7.4.3

(A) DC output characteristics of a BGFET device at 250°C; (B) DC output characteristics of a TGFET device at 200°C [7].

achieved for all BGFET devices due to the isolated back gate. However, at 250°C, a DC output characteristics as presented in Fig. 2.7.4.3A is observed. A maximum drain current density of more than 100 mA/mm at a drain bias of <20 V was extracted with a maximum extrinsic transconductance of 7.5 mS/mm. Most of the TGFET devices could be pinched off at room temperature, although the current levels were small due to incomplete boron activation. The output characteristics of a TGFET device at an operation temperature of 200°C are presented in Fig. 2.7.4.3B. The gate length and width for this device were 20 and 100 μm, respectively. The maximum extrinsic transconductance was 1.4 mS/mm. At a gate length of 10 mm, the maximum current density was 40 mA/mm; however, the device could not be pinched off completely. This indicates that the work mode of TGFET was affected by the gate length. It should be noted that no clear off-current state was shown in the two devices. Kato et al. proposed a skillful technique of selective phosphorus-doped n-type diamond growth using crystalline orientation dependence of growth rate and phosphorus incorporation [8, 9]. This concept was applied by Iwasaki et al. [10] to control depletion layers in diamond JFETs with lateral pn junctions, and clear off-states were observed. The device could be operated in both depletion and enhancement modes. Moreover, the vertical device structure is preferred for power device applications [11–13]; JFETs consisting of lateral pn junctions are important to be developed. Fig. 2.7.4.4A and B illustrate the simplified top-view and cross-sectional schematics of the lateral JFET structure, respectively. A bar-shaped p-channel was sandwiched by two n+-diamond regions to form lateral pn junctions at the interfaces. The depletion layers in the p-channel were controlled by gate voltages. A difference of three orders of magnitude in the dopant concentrations in the n and p diamond regions was achieved to guarantee the depletion layers extending mainly into the p-channel, modulating the drain current passing through the p-channel. The buffer diamond layer was inserted to obtain a smooth p-channel layer. Fig. 2.7.4.4C and D show secondary electron microscopy (SEM) images of a fabricated JFET with a channel width of 1 μm. The p-channel was observed at the center indicated by the two dashed lines in Fig. 2.7.4.4D. The inclined regions at both sides of the p-channel are selectively grown n+-diamonds.

236

2. Semiconductor diamond

FIG. 2.7.4.4

Diamond JFET structure. Schematics in (A) top view and in (B) cross-section along the dashed line in panel (A). (C) Top-view SEM image of a 1-m device. (D) Magnified SEM image of the region indicated by the white box in panel (C). The lattice orientation shown in panel (A) is also applicable to panels (C) and (D) [10].

The JFETs with lateral pn junctions operated successfully as a switching device by controlling the depletion layer at the interface of the lateral pn junctions, as presented in Fig. 2.7.4.5A and B. The devices could be in the off state at a gate voltage over 6.7 and 30 V for the 0.5 and 1 μm devices, respectively. Additionally, high current on/off ratios in the range of 107–108 and steep subthreshold swings of 95–120 mV/decade can be obtained. Because of the large bandgap and high thermal conductivity, diamond devices are expected to work at high temperatures [14, 15]. Therefore, the lateral JFETs also can work at a high temperature [16]. Fig. 2.7.4.6A and B show the IDS-VDS curves of a diamond JFET measured at 300 and 673 K, respectively. The linear and saturation regions were clearly observed for both temperatures. A high current density of 1300 A/cm2 (3.2 mA/mm) was obtained at 673 K, 50 times higher than 25 A/cm2 (0.06 mA/mm) at 300 K. Here, the drain current densities in A/cm2 and in mA/mm were calculated by normalizing with the channel cross-section and thickness, respectively. For JFETs, the unipolar current modulation can be kept until the gate bias reaches the builtin potential of the p-n junction between the gate and channel. The ideal built-in potential of the diamond p-n junction is 4.5 V. The built-in potential was reduced by 0.34 V at 673 K, according to calculations. Thus, the modulation of the drain current saturated at lower gate

2.7.4 Junction FETs

237

FIG. 2.7.4.5 Id-Vd curves of JFETs with different channel widths of (A) 0.5 and (B) 1 μm. The measurements were performed at room temperature. The channel thickness was 0.7 μm [10].

FIG. 2.7.4.6 IDS-VDS curves of diamond JFET. (A) 300 K. (B) 673 K [16].

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2. Semiconductor diamond

voltages at higher temperatures. The very low leakage current around 1014 A was maintained at high temperatures up to 723 K. This is the advantage of the wide bandgap of diamond. Because diamond JFETs showed high on currents and low leakage, high on/ off ratios over 106 in the whole temperature range were obtained. This increased with temperature elevation and reached the maximum of 4 107 at 623 K. The breakdown of the lateral JFET occurred at the p-n junctions between the p-channel and n+-gates [17]. At room temperature, when the source-drain bias was 566 V with a gate bias of 20 V, the device was broken down, as shown in Fig. 2.7.4.7A. The gate-drain p-n junction sustained a higher voltage of 586 V at this point, corresponding to a high electric field of 6.2 MV/cm, higher than the physical limitation of 4H-SiC and GaN [18]. The high electric field resulted from the high-quality p-n junctions formed by the selective growth of the n+-diamonds under the optimized MPCVD conditions. Note that the value was estimated with the assumption that the gatedrain junction was composed of a planar one-side abrupt p-n+ junction. When the temperature was increased to 200°C, a higher breakdown voltage of 608 V was achieved, as shown in Fig. 2.7.4.7B. Observation of the higher breakdown voltage at a higher temperature is a strong indication of the avalanche breakdown [19]. Acceleration of minority carriers in the reverse biased p-n junction is suppressed at a higher temperature because of the increase of phonon scattering, resulting in a decrease in the impact ionization rate. Thus, diamond JFETs with good blocking characteristics can be fabricated by the selective growth method. Normally off diamond JFETs were also demonstrated; these can have higher threshold voltages than those from other semiconductors such as Si and SiC because diamond p-n junctions possess a higher built-in potential of 4.5 V. Therefore, a diamond semiconductor has a large voltage margin between the threshold voltage and the leakage of the gate diodes. Due to a low dielectric constant and a high doping concentration, a narrow channel width is necessary for the fabrication of normally off diamond JFETs. Suwa et al. fabricated normally off diamond JFETs with 0.26 μm channel width, and those SEM images are presented in Fig. 2.7.4.8 [20]. Fig. 2.7.4.9 shows the output characteristics of the normally off diamond JFET at room temperature. The I-V curves were measured by changing the gate voltage from 4 to 0 V. Without applying the gate voltage, no drain current was observed, indicating the normally off operation of the diamond JFET. Even though the boron concentration of the channel increased

–4.0

–1.5

1.0

–7

1.0–9

1.0–11

–1.0

0

–200

–400

–600

–1.0

Drain current (A)

1.0–5

Drain current (µA)

–2.0

200°C Vg = 20 V

566 V

–3.0 Drain current (A)

Drain current (µA)

RT Vg = 20 V

–0.5

(A)

0

–200 –400 –600 Drain voltage (V)

1.0–5 1.0–7 1.0–9

1.0–11 0

Drain voltage (V)

0

608 V

–200

–400

–600

Drain voltage (V)

0

(B)

0

–200 –400 –600 Drain voltage (V)

FIG. 2.7.4.7 Blocking characteristics of diamond JFET at (A) RT and (B) 200°C. The insets in the panels show leakage currents in a logarithmic scale [17].

References

239

FIG. 2.7.4.8 SEM images of the diamond JFET (A) after ICP etching and (B) after n+ selective growth [20].

FIG. 2.7.4.9 Output characteristics of diamond JFET with a channel width of 0.26 μm [20].

gradually to the surface, the channel was depleted owing to the tapered channel. A large threshold voltage (3.0 V) was in good agreement with the calculations. In fact, the threshold voltage could be controlled by the channel width and dopant concentration. In conclusion, research on the JFETs was mainly focused on the selective growth method. The threshold voltage was well controlled by the channel width and dopant concentration. When a proper channel width is applied, diamond JFET can enter the normally off mode, indicating a promising application in power devices.

References [1] T.H. Borst, O. Weis, Electrical characterization of homoepitaxial diamond films doped with B, P, Li and Na during crystal growth, Diam. Relat. Mater. 4 (1995) 948. [2] K. Okano, H. Naruki, Y. Akiba, T. Kurosu, M. Iida, Y. Hirose, T. Nakamura, Characterization of boron-doped diamond film, Jpn. J. Appl. Phys. 28 (1989) 1066. [3] A. Vescan, I. Daumiller, P. Gluche, W. Ebert, E. Kohn, High temperature, high voltage operation of diamond Schottky diode, Diam. Relat. Mater. 7 (1998) 581. [4] J.E. Field, The Properties of Diamond, Academic Press, London, 1979. [5] T.H. Borst, S. Strobel, O. Weis, High-temperature diamond p-n junction: B-doped homoepitaxial layer on Ndoped substrate, Appl. Phys. Lett. 76 (1995) 2651.

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2. Semiconductor diamond

[6] Y. Gurbuz, O. Esame, I. Tekin, W.P. Kang, J.L. Davidson, Diamond semiconductor technology for RF device applications, Solid State Electron. 49 (2005) 1055–1070. [7] A. Aleksov, A. Vescan, M. Kunze, P. Gluche, W. Ebert, E. Kohn, A. Bergmaier, G. Dollinger, Diamond junction FETs based on δ-doped channels, Diam. Relat. Mater. 8 (1999) 941–945. [8] H. Kato, T. Makino, M. Ogura, N. Tokuda, H. Okushi, S. Yamasaki, Selective Growth of Buried n+ Diamond on (001) Phosphorus-Doped n-Type Diamond Film, Appl. Phys. Express 2 (2009) 055502. [9] H. Kato, T. Makino, M. Ogura, D. Takeuchi, S. Yamasaki, to be published in Japan. Maskless Selective Growth Method for p-n Junction Applications on (001)-Oriented Diamond, Jap. J. Appl. Phys. 51 (2012) 090118 [10] T. Iwasaki, Y. Hoshino, K. Tsuzuki, H. Kato, T. Makino, M. Ogura, D. Takeuchi, T. Matsumoto, H. Okushi, S. Yamasaki, M. Hatano, Diamond Junction Field-Effect Transistors with Selectively Grown n+ -Side Gates, Appl. Phys. Express 5 (2012) 091301. [11] K. Ikeda, H. Umezawa, K. Ramanujam, S. Shikata, Thermally Stable Schottky Barrier Diode by Ru/Diamond, Appl. Phys. Express 2 (2009) 011202. [12] H. Umezawa, K. Ikeda, R. Kumaresan, N. Tatsumi, S. Shikata, Increase in reverse operation limit by barrier height control of diamond Schottky barrier diode, IEEE Electron Device Lett. 30 (2009) 960. [13] K. Tone, J.H. Zhao, L. Fursin, P. Alexandrov, M. Weiner, 4H-SiC normally-off vertical junction field-effect transistor with high current density, IEEE Electron Device Lett. 24 (2003) 463. [14] A. Vescan, I. Daumiller, P. Gluche, Very high temperature operation of diamond Schottky diode, IEEE Electron Device Lett. 18 (1997) 556–558. [15] H. Umezawa, Y. Kato, S. Shikata, 1Ω on-resistance diamond vertical-Schottky barrier diode operated at 250°C, Appl. Phys. Express 6 (2012) 011302-1–011302-4. [16] T. Iwasaki, Y. Hoshino, K. Tsuzuki, H. Kato, T. Makino, M. Ogura, D. Takeuchi, H. Okushi, S. Yamasaki, M. Hatano, High-temperature operation of diamond junction field-effect transistors with lateral p-n junctions, IEEE Electron Device Lett. 34 (9) (2013). [17] T. Iwasaki, J. Yaita, H. Kato, T. Makino, M. Ogura, D. Takeuchi, H. Okushi, S. Yamasaki, M. Hatano, 600 V Diamond junction field-effect transistors operated at 200°C, IEEE Electron Device Lett. 35 (2014) 2. [18] A. Hiraiwa, H. Kawarada, Figure of merit of diamond power devices based on accurately estimated impact ionization processes, J. Appl. Phys. 114 (2013) 034506-1–034506-9. [19] Y. Lee, M. Han, Y. Choi, Analytic models for the temperature dependence of the breakdown voltage of 6H- and 4H-SiC rectifiers, J. Korean Phys. Soc. 39 (2001) 20–22. [20] T. Suwa, T. Iwasaki, K. Sato, H. Kato, T. Makino, M. Ogura, D. Takeuchi, S. Yamasaki, M. Hatano, Normally-off diamond junction field-effect transistors with submicrometer channel, IEEE Electron Device Lett. 37 (2) (2016).

2.7.5.1 Introduction

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2.7.5 Solar-blind deep-ultraviolet detectors Meiyong Liaoa, Xiaohui Changb, Zhangcheng Liub, Hongxing Wangb, Yasuo Koidea a

Research Center for Functional Materials, National Institute for Materials Science (NIMS), Tsukuba, Japan bInstitute of Wide Band Gap Semiconductors, Xi’an Jiaotong University, Xi’an, P.R. China

2.7.5.1 Introduction A photoelectric detector is a device transferring the optical energy to the electrical signal. Deep ultraviolet (DUV: 350–190 nm) has drawn much attention for various applications in environmental security, information technology, medical treatment, astronomical observation, the military, and intersatellite communications. The sun is a natural UV source of which the stratospheric ozone layer absorbs the light with a wavelength shorter than 280 nm. Therefore, the UV light reaching the Earth’s surface is in the range of 280–400 nm, which is usually divided by UVA (400–320 nm) and UVB (320–280 nm). For DUV detection, it is desirable to utilized solar-blind photodetectors, which do not or response weekly to the light with wavelength larger than 280 nm. A traditional UV-enhanced Si photodetector has some limitations in UV detection, owing to its narrow bandgap energy of 1.1 eV. Therefore, the Si photodiode suffers from high temperature instability for DUV detection [1]. Thus, researchers began to develop photodetectors on wide bandgap semiconductors (WBG) such as GaN [2], ZnO [3], SiC [4], and Ga2O3 [5]. Diamond has also become an extraordinary candidate for ultraviolet photodetectors thanks to its wide bandgap, high carrier mobility, radiation hardness, and thermal stability. When using these semiconductors for DUV detection, filters have to be added to reduce the background noise from the solar light. True solar blindness (cut-off wavelength <280 nm) can be achieved by using diamond for photon detection due to its wide bandgap of 5.5 eV. The wide bandgap energy of diamond also offers the advantage of an extremely low dark current. An ideal diamond DUV detector should satisfy the 5S requirement of Sensitivity, high Signal-to-Noise ratio, high spectral Selectivity, high Speed, and high Stability [6, 7]. The photoconductive behavior of diamond can be traced back to 1934, when Robertson measured the photocurrent of different natural diamonds [8]. Based on the optical absorption critical energy, diamond was classified into Type I and Type II. The transient photocurrent behavior was also investigated in the 1950s to deduce the carrier lifetime and mobility [9, 10]. Nevertheless, the investigation into

242

2. Semiconductor diamond

diamond as a photodetector began in the 1990s, first on polycrystalline diamonds [11–14] and then on single-crystal diamonds. In recent years, owing to the development of the microwave plasma chemical vapor deposition (MPCVD) technique, high-quality single-crystal diamonds (SCD) have been developed. SCD is desirable to achieve the above 5S requirements due to the nonexistence of grain boundaries and nondiamond impurities that exist in polycrystalline diamonds (PCDs). Photodetectors with different structures based on SCDs have been examined. However, compared to PCDs, SCDs have some limitations such as a small size and high cost. Thus, UV detectors based on PCDs have also been explored. In this chapter, we mainly describe SCD photodetectors with a brief introduction of PCD ones.

2.7.5.1.1 Basic parameters Responsivity, quantum efficiency, and gain: Responsivity, Rλ, is described as the output signal per radiant energy, P, in watts. In most cases, photocurrent in amperes, Ip, is directly measured as the output. The responsivity is expressed as Rλ ¼

Ip P

(2.7.5.1)

Note that the energy is usually expressed as W/cm2 and the photocurrent is A/cm2. The radiant energy P is the product of the photon energy hν and flux F. P ¼ F･hν

(2.7.5.2)

By inducing the quantum efficiency η, the generation rate of electrons can be obtained as G ¼ η･F

(2.7.5.3)

Therefore, the responsivity can be rewritten as Rλ ¼ η

λðnmÞ   W 1240 nm  A

(2.7.5.4)

where the unit of the light wavelength λ is nm. As a general, one can add gain, g, in the expression of Rλ as Rλ ¼ η･g･

λðnmÞ   W 1240 nm  A

(2.7.5.5)

There is another similar parameter of Sensitivity, which is the change of output with the change of input light. Spectral selectivity: A perfect photodetector absorbs only the light near one certain wavelength band and is transparent to all other bands. Experimentally, one can obtain the spectral selectivity by plotting the photocurrent as a variation of the wavelength of the incident light.

2.7.5.2 Device structures

243

The spectral selectivity can be simply achieved by selecting the bandgap energy of the semiconductor. A true solar/visible-blind DUV detector requires a high rejection ratio of the photocurrents between the DUV light and those lights with a wavelength no longer than 280 or 400 nm. Dark current: Dark current denotes the electrical current under an external bias without illumination. It is related to the electric conductivity of the semiconductor, which is determined by the bandgap of the semiconductor, the dopants, and the temperature. Dark current contributes to the noise of the photodetector. Response time: The response time refers to the transient behavior of the current with switching on/off the incident light. Therefore, both the rise time and fall time can be described. Technically, one can define rise time (Tr) as the output level changing from 10% to 90% of the peak output level. The fall time (Tf) corresponds to the time required for the photodetector output level decreasing from 90% to 10% of the peak output level. In most cases, the response time of a photodetector is determined by the defects. It can also be limited by the capacitive effects and the device geometries. Noise equivalent power and specific detectivity: Noise equivalent power (NEP) defines the sensitivity of a photodetector, which is equal to the noisepspectral density ffiffiffiffiffiffi pffiffiffiffiffiffi divided by the responsivity. Thus, the unit of NEP is expressed as A/ Hz or V/ Hz. The specific detectivity (D*) is defined as the reciprocal of NEP normalized per square root of the sensor’s area and frequency bandwidth.

2.7.5.2 Device structures Different device structures can be adopted for the photodetectors, as shown in Fig. 2.7.5.1. The most-used geometry is the metal-semiconductor-metal (MSM) structure, which offers the simplest device fabrication. In the MSM structure, the electrical contacts can be either Schottky or ohmic, usually with an interdigitated finger arrangement. When symmetrical electrodes are used, no photocurrent can flow through the devices at zero bias. The MSM photodetector with two back-to-back Schottky contacts offers the advantage of low dark current and high response speed, but the quantum efficiency is less than unity. The MSM structure with two ohmic contacts is called a photoconductor, which can show a very high external quantum efficiency much larger than unity. One drawback of the MSM photodetector is the exposure of the optical receiving area, which may suffer from surface pollutants. Another conventional device structure is a Schottky photodiode (SPD) with ohmic and Schottky contacts. There are several advantages for the SPD: high speed, zero-voltage operation, and protection of the optical receiving area by the Schottky contact. However, the Schottky contact needs to be semitransparent, which induces incident optical loss. The pn-junction type structure, the basic structure for the solar cell, is also often used as a photodiode. This photodiode can be operated at zero bias or in photovoltaic mode. Most of the photodetector structures mentioned above were reported on SCDs. The phototransistor structure is less reported due to the complex fabrication process. For diamond, due to the difficulty in n-type doping, a PNP phototransistor has not been reported yet.

244

FIG. 2.7.5.1

2. Semiconductor diamond

Basic device structures for photodetectors.

2.7.5.3 Detectors based on polycrystalline diamond Polycrystalline diamonds have the advantage of being able to grow on a silicon substrate for large-scale production. However, the photoconductivity of PCDs was restricted by the grain defects, impurities, and density of the grain boundary, which were directly affected by the grain size. Tang et al. evaluated the effect of different diamond grain sizes on the performance of the photodetector, and indicated that the dark current was reduced with increasing grain size, and the ratio of photocurrent/dark current and the net photocurrent had been improved for the larger grain size [15, 16]. The photodetector based on PCDs usually exhibited persistent photoconductivity (PPC), which greatly extended the response time, not suitable for detecting high frequency signals. The PPC phenomenon mainly originates from minority carrier trapping. The dynamic behaviors of the minority carriers in microcrystalline diamond films and PPC were investigated by charge-based deep-level transient spectroscopy [17, 18]. It was shown that 23 min were needed for the photocurrent to decay to 10% of the saturated value, followed by more time for the photocurrent to get back to the dark current [18]. The charge-based deep-level transient spectroscopy technique was used to examine the trapping centers in diamond. The spectra were measured by deep-level transient spectroscopy (DLTS) over a temperature range of 255–289 K. It was found that a shallow energy level with activation energy Ea ¼ 0.213 eV existed within the bandgap, contributing to the PPC. In addition, enhancement of photoconductive gain and responsivity were primarily due to the trapping effect of the shallow level, which could reduce the carrier recombination probability [19]. Polycrystalline diamond UV photodetectors could not reach ideal performance due to the fact that carriers were scattered by the grain boundaries during the transport process. Wei et al. fabricated a graphene/microcrystalline diamond (MCD) heterojunction vertical structure UV detector to avoid the influence of the grain boundary on the transportation of carriers [19]. In their device structure, a monolayer was graphene was transferred on top of undoped and boron-doped microcrystalline diamond films as top electrodes, and the metal

2.7.5.4 Detectors on single-crystalline diamonds

245

films of Ti/Au were deposited and annealed at the bottom of the diamond as another electrode. A contact pad made of SiO2/Ti/Au was deposited on the top surface of the microcrystalline film. The responsivity for the graphene/pMCD device reached 1.4 A/W, which was higher than those of the graphene/MCD device and various Schottky UV photodetectors. The high responsivity of the graphene/pMCD device was attributed to the following facts: (1) the transport direction of the carrier was parallel to the grain boundary in the device with a vertical structure, thereby reducing the losses caused by grain boundary scattering; (2) monolayer graphene as a transparent electrode could be penetrated by 98% ultraviolet light, leading to the photosensitive area increasing and device performance improving; (3) The defects at the graphene/p-MCD interface might lead to a photocurrent gain. The defects at the metal/semiconductor interface can trap photogenerated carriers and induce a reduction of the Schottky barrier height upon illumination, leading to a photocurrent gain [11]. In addition, the photoresponse properties of photodetectors fabricated from PCDs could be markedly enhanced by a surface treatment [11]. Nanocrystalline diamond photodetectors were also reported. Lin et al. investigated UV photodetectors fabricated by nanocrystalline diamond film that had a high sp3 fraction and a low surface roughness [20]. Au interdigital electrodes were utilized as the Ohmic contact to nanocrystalline diamond by annealing at 500°C in vacuum. It was observed that the dark current was as high as 0.2 mA, although the rise and fall time were 0.8 and 1.2 s, respectively.

2.7.5.4 Detectors on single-crystalline diamonds Single-crystal diamond is desirable to develop DUV detectors to satisfy the 5S requirements. In recent decades, large-area, low-cost SCDs have become available with the development of CVD technology. By virtue of the great progress in SCD crystal growth, DUV detectors with high performance have been reported. Symmetrical MSM photodetectors: For DUV light with a wavelength less than 225 nm close to the bandgap energy of diamond, most of the light will be absorbed near the surface region of diamond. However, for a longer wavelength, the light will penetrate through the intrinsic diamond layer with little absorption. In most cases, homoepitaxial SCD layers were deposited on a high-pressure high-temperature (HPHT) substrate or a CVD SCD substrate for DUV photodetector fabrication. The type-Ib diamond substrates, which are the most used, contain a large amount of nitrogen impurities and sometimes can have dislocations. Therefore, the photo carriers generated within the diamond bandgap can also be collected even when the photon energy is less than the bandgap energy. A preliminary MSM photodetector with a Ti contact was fabricated on a 20-μm-thick SCD epilayer by Teraji et al. [21]. The solar-blind ratio, defined as the photocurrent rejection ratio between the 220–280 nm light, was around three orders of magnitude. The thick high-quality SCD layer should achieve high-performance diamond DUV detectors. However, the quantum efficiency and the transient response were not provided in that work. By properly designing the device structures, a thick epilayer layer is not always necessary to achieve high responsivity and a solar/visible blind ratio. Liao et al. reported a planar interdigitated-finger

246

2. Semiconductor diamond

type MSM photodiode fabricated from a submicron-thick SCD epilayer grown on a type-Ib HPHT diamond substrate [22]. Due to the p-type nature of the epilayer and the n--type nature of the diamond substrate, a subsurface p-n junction was formed between the epilayer and the substrate, which depleted the holes in the epilayer. For a 500-nm boron-doped diamond epilayer, when [B] is around 1015cm3, all the holes can be depleted without any external bias. Therefore, by using the subsurface p-n junction, the photoresponse properties such as the dark current, responsivity, the solar/visible-blind ratio, and the response speed of the diamond DUV photodetector can be tailored through boron-doping engineering. The device contacts were composed of interdigital-finger (IDF) Ti/WC electrodes. The Ti/WC contact as above was sputter deposited on the oxidized diamond epilayer as the metal contacts. The finger spacing and the device active area were 10 μm and 52 103 mm2, respectively. The as-fabricated MSM photodetector was an MSM photodiode (PD) with a metal/semiconductor Schottky contact. The Schottky barrier height ranges from 1.2 to 2.1 eV due to the Fermi-level pinning of the oxygen-terminated diamond surface [23]. The MSM PD turned out to be a photoconductor with Ti/WC ohmic contacts after annealing at 600oC. The MSM PD showed a very low dark current beyond the semiconductor analyzer of 0.1 pA due to the existence of the interface barrier. Due to the depletion of the holes in the p-type epilayer, the photoconductor fabricated on an oxidized epilayer also had a low dark current (1 pA), even at a bias of 20 V. The dark current-voltage (I-V) curve disclosed a space-charge-limited-current (SCLC) behavior, consistent with ohmic contacts to a highly resistive material [24], as shown in Fig. 2.7.5.2A. Upon 220-nm DUV light illumination, a marked difference in the I-V characteristics was observed for the MSM-PD and photoconductor, as revealed in Fig. 2.7.5.2B. The photocurrent of the MSM-PD increased as the voltage increased due to the extension of the depletion layer width and the modification of the interface states. For MSM-PD fabricated on the epilayer based on the subsurface pn-junction, no photoconductivity gain was observed (or the quantum efficiency <1) in the voltage range of 0–30 V for the photodiode, even for a 220-nm light intensity of 20 μW/cm2. In contrast, the MSM photoconductor using the same Ti/WC contact after annealing exhibited a much larger photocurrent than the photodiode. The photocurrent of the photoconductor at 2 V was about four orders of magnitude larger than that of the MSM-PD and tended to saturate with increasing the bias. The saturation behavior of the photoconductor can be interpreted by the space-charge-limited photocurrent. A responsivity of 6 A/W at 220 nm was obtained for the photoconductor at a bias of 3 V, indicating a gain of 33, if assuming the external quantum efficiency unity. The mobility-lifetime product μτ was calculated to be around 105 cm2 V1, taken from the linear regime in the I-V curve to avoid the space-charge-limited-current effect. g¼

μτV L2

(2.7.5.6)

where L is the gap between the electrodes. Note that a uniform electric field distribution was assumed for the calculation. The spectral response revealed that the highest responsivity was observed at 210 nm, as shown in Fig. 2.7.5.2C. A shoulder at around 270 nm also appeared, independent of the device structure of a photodiode or a photoconductor. The 270-nm response is due to the type-Ib substrate effect, which contains a large amount of nitrogen. Rohrer et al. studied the absorption of diamond films with different nitrogen content and found that the intensity was

2.7.5.4 Detectors on single-crystalline diamonds

Dark current (A)

FIG. 2.7.5.2

(A) Dark I-V characteristics of the MSM photoconductor. (B) I-V characteristics of the photoconductor and MSM-PD during 220 nm light illumination. (C) Spectral photoresponse measured from the MSM-PD and the photoconductor. The showed responsivity indicates a gain for the photoconductor [22].

10–12

10–13

10–14

0

5

(A)

10 15 Applied voltage (V)

20

10–7 10

101

Photocurrent (A)

100 10–9 10–1 10–10 10–2

10–11

MSM-PD 10–3

10–12

10–4 –13

(B)

Responsivity (a.u.)

10–1

0

5

10 15 Applied voltage (V)

20

MSM PD MSM Photoconductor

10–3

10–5 Subband gap 10–7

10–9 200 250 300 350 400 450 500 550 600 650 (C) Wavelength (nm)

Responsivity (A/W)

MSM Photoconductor

–8

10

247

248

2. Semiconductor diamond

strongly related to paramagnetic nitrogen density [25]. At 20 V, the discrimination ratio between the 210 nm and the visible light (400 nm) was five orders of magnitude for the MSMPD. It reached nearly 108 for the photoconductor, which is still the largest value in diamond photodetectors up to now. A weak photoresponse at the wavelength larger than 500 nm was observed, which may be associated with boron or nitrogen. This subband signal disappeared after preilluminating the device with white light. The spectrum shapes of the photoconductor and MSM-PD were similar. Both the MSM PD and the photoconductor exhibited quick response. Fig. 2.7.5.3A shows the time response upon the 220-nm light (intensity 20 μw/cm2) illumination measured by a dc picoammeter using the mechanical way. The predominant response time at the rising and falling stages was beyond the measurement system time constant of 0.3 s while a slow component with small amplitude appeared during the decay. The possible origins were FIG. 2.7.5.3 (A) Time response of the

10–7

Photocurrent (A)

10–8 10–9

Slow component

10–10 10–11 10–12 10–13 10–14

0

100

200

(A)

300 Time (s)

400

500

3.0 2.5 Amplitude (V)

photoconductor upon the 220-nm light illumination measured by a mechanical chopping method, and (B) time-resolved photoresponse upon a 193-nm excimer laser recorded by an oscilloscope with a 50-Ω impedance [22].

2.0 1.5 1.0 0.5 0.0 –0.5 –150

(B)

–100

–50

0 Time (ns)

50

100

150

2.7.5.4 Detectors on single-crystalline diamonds

249

boron-related impurities or interface traps. The weak persistent photoconductivity (PPC) indicated that the density of the trap states was rather low. The response to a 193-nm excimer laser was also measured, which is shown in Fig. 2.7.5.3B. The full width at half maximum (FWHM) of the photoresponse was around 10 ns, which was almost the same value of the laser pulse width. No dependence of the FWHM on the bias was observed. When an H-terminated diamond surface was used for the MSM device application, a huge DUV photocurrent gain was observed, with the highest responsivity of 230 A/W [26]. The 100 nm-thick boron-doped diamond on the HPHT type-Ib diamond forms the subsurface junction, depleting the holes in the epilayer. Therefore, by combination with the Ti Schottky contact on the H-diamond surface, the MSM structures showed a very low dark current of pA in vacuum and 100 pA in air. Schottky photodiode: A typical Schottky photodiode (SPD) contains a semitransparent Schotty contact and an Ohmic contact. Here, we call this device the conventional SPD (CSPD) to distinguish the one described later [27, 28]. The CPSD has the advantage of the simultaneous passivation effect of the optical receiving area by the Schottky contact. The CSPD was fabricated on the SCD epilayer with a boron concentration of around 1016 cm3. The ohmic contact was a Ti layer with a thickness of 40 nm followed by a WC protection cap layer with a thickness of 30 nm, annealed at 600oC for 1 h in an argon ambient to achieve a good ohmic property. A semitransparent WC Schottky contact was deposited on a defined circle pattern with a diameter of 400 μm. The structure of the resultant WC thin films was cubic with (111) preferential orientation [29]. The electrical resistivity and the composition of the WC thin film changed little after annealing at 500oC. The optical transmittance of the WC thin films was evaluated to be about 70% in the visible light region and 58% at 220 nm for a 3-nm thickness. The as-fabricated diamond CSPD showed typical rectifying properties with a rectification ratio of 108 at 5 V in dark condition. The n and qΦB values of the as-fabricated device were evaluated to be 1.67 and 1.59 eV, respectively, deduced from the thermionic emission theory [30]. Transition carbides or nitrides are noncarbide-forming materials, which are promising candidates as Schottky contacts to diamond due to the metallic conductivity, chemical inertness, and oxidation resistance. It was shown that the WC Schottky contact had thermal stability up to 500oC, even annealing for 5 h. The dark current kept less than 0.1 pA, the SBH was larger than 1 eV, and the n value was 1 after annealing at 500oC. The as-fabricated CSPD showed a low photocurrent upon the 220-nm light illumination at reverse biases. A saturation was observed when the bias was larger than 5 V. The saturation photocurrent, Iphs, was approximately 1010 A under the light intensity around 20 μW/cm2. The as-fabricated SPD had a low quantum efficiency η of 2% (assuming g ¼ 1) and R of 4 mA/W at a reverse bias larger than 5 V without correction for surface reflection. After annealing for 1 h, Iphs increased by a factor of 2. However, a marked enhancement of the photocurrent was observed after annealing for 2 h or longer. Iphs reached 107 A, which was enhanced by a factor of 103 compared with the as-fabricated photodiode. The photodiode after annealing provided a gain larger than unity. No deterioration of the enhanced photocurrent occurred upon annealing at 500oC for 5 h. At 220-nm light with an intensity of 20 μW/cm2, the rise time Tr was 10 s and fall time Tf was 300 s for the as-fabricated SPD. After annealing for more than 2 h, PPC was clearly observed, although the rise time changed little. By using the merits of SPD in photovoltaic mode, the PPC effect can be overcome [31]. In this mode, the photo-generated electron-hole pairs are driven by the internal built-in field

250

2. Semiconductor diamond

FIG. 2.7.5.4 (A) Time response of SPD and (B) linearity of the SPD DUV sensor at zero bias [31].

in the depletion region, forming the current when the device is short-circuited. The carrier injection from the electrodes was avoided at zero bias, which resulted in a quick response time. Fig. 2.7.5.4A presents the time response at zero bias. The result disclosed that the rising and falling times were beyond the limit of 0.3 s of the system. The response time did not depend on the incident light intensity. When compared with the reverse-bias mode, the response time and linearity were significantly improved in the photovoltaic mode (Fig. 2.7.5.4B). A responsivity of 0.5 mA/W at 220 nm was obtained at zero bias. Interdigital-finger SPD: The electrical contacts of the MSM photodetectors are usually symmetrical and operated in one mode. On the other hand, the CSPD are constructed from an ohmic contact and a semitransparent circled Schottky contact. The CSPD has the drawback of photon absorption in the Schottky metal, leading to the reduction of quantum efficiency and the presence of a subbandgap photoresponse due to internal photoemission from the metal to the semiconductor. For the sake of high photo responsivity and high response speed, SPDs with IDF ohmic and Schottky contacts were created, as shown in Fig. 2.7.5.5A [32]. The IDF-SPD combined the advantages of the MSM PD and the photoconductor. The ultrawide bandgap of diamond and the subsurface junction offer the opportunity to fabricate the IDF-SPD with low dark current at both the forward and reverse biases. The IDFSPD could have high photoresponsivity as a photoconductor at forward biases due to the ohmic contact nature and the high speed at reverse bias due to the Schottky contact nature. By controlling the boron concentration in the epilayer to less than 1016 cm3 and the thickness

2.7.5.4 Detectors on single-crystalline diamonds

251

FIG. 2.7.5.5 (A) Device geometries of the IDF-SPD and (B) dependence of photocurrent and responsitivity of the IDF-SPD on the applied bias upon the 220-nm light illumination [32].

of the epilayer to less than 500 nm, the dark current of the IDF-SPD was lower than 1014 A at reverse biases and around 1012 A at forward biases up to 30 V. Fig. 2.7.5.5B shows the I-V characteristics of the IDF-SPD upon illuminating the device with the 220-nm light, where the right vertical axis is the corresponding responsivity. The I-V curve clearly displayed a rectifying nature with a rectification ratio of more than 104 at 5 V during illumination. The responsivity dependence on the reverse bias was similar to the symmetrical MSM-SPD but different from the CSPD, where the photocurrent was saturated at low biases. The external quantum efficiency reached 20% at a reverse bias of 30 V and was larger than 100% for VF <1 V at forward biases. The transient response upon the 220-nm light chopped at 100 Hz at a forward bias of 32 V revealed a dominated response time larger than 10 ms. In contrast, the response time was smaller than 10 ms at a reverse bias of 32 V because the photoresponse followed the 100 Hz chopped light pulse repeatedly. The response time was much faster in the reverse-bias mode than in the forward-bias mode. The spectral response of the IDF-SPD was similar to that of the MSM-PD and photoconductor. The IDF-SPD makes use of the advantages of the traditional interdigitated MSM device in charge collection. It behaves as an MSM photodiode at reverse biases with high speed and an MSM photoconductor with gain at forward biases. The maximum gain-bandwidth products in both modes are similar due to the same electrode distance. When compared with the CSPD, the responsitivity of the IDF-SPD was improved. The photocurrent in a CSPD saturates at low reverse biases with a typical low efficiency, no more than one-fourth that of the IDF-SPD. The results above reveal that a thin diamond film is enough to achieve a large visible-blind ratio by using the IDF structures (MSM or IDF-SPD) where nontransparent electrodes are used. The Ib (100) diamond substrate contributes little to the spectral response in the visible-light regime, even at high electric fields for the MSM or IDF-SPD device structures. This is quite different from the CSPD where the absorption occurs mainly beneath the semitransparent Schottky contact, leading to visible light response at even smaller reverse biases.

252

2. Semiconductor diamond

2.7.5.5 Tailoring the photoresponse properties The photoresponse properties depend on the types of SCD, namely, the impurities, the epilayer thickness, and the surface states in addition to the device design discussed above. Diamond types: Photodetectors fabricated on the commonly available SCD substrates have shown poor photoresponse properties. The spectral response from type-Ib diamond with an oxygen-terminated surface is governed by nitrogen absorption and the DUV photocurrent is rather low with a typical quantum efficiency of 0.1%. Natural type-IIa diamonds also have uncontrollable defects. Fig. 2.7.5.6A illustrates the spectral responses of the photodetector directly fabricated on the type Ib and natural type II a SCD substrates, which are far from ideal. Therefore, photodetectors have been usually developed by growing high-quality homoepitaxal diamond layers on SCD substrates. The spectral response of the diamond DUV detector can be tailored by the epilayer thickness, as shown in Fig. 2.7.5.6B, which was normalized by the incident photon flux. The sharp band-edge response is clearly seen for all the photodetectors. For the submicron thin diamond epilayer grown on the type Ib diamond substrate, the shoulder at around 270 nm is observed. A visible light response at 450–600 nm was also observed. As the films became thicker, the 270-nm photoresponse turns out to be less distinct and finally disappears for the for the 5-μm-thick epilayer. The visible light absorption from 630 to 450 nm likely originates from the Ib diamond substrate. Surface terminations: Oxygen-terminated diamond (O-diamond) and hydrogen-terminated (H-diamond) surfaces are the most utilized for diamond DUV detectors. In most cases, O-diamond is desirable to achieve an extremely low dark current. Unless the diamond used for device fabrication is properly designed, photodetectors fabricated from H-terminated diamond always exhibit high dark current due to the existence of 2DHG surface conductivity. The photocurrent is buried in the dark current for the photodetector fabricated on an asgrown 500-nm undoped H-diamond epilayer on the type Ib HPHT diamond substrate

FIG. 2.7.5.6

(A) Spectral response of the photodetectors fabricated directly on the type Ib and IIa diamond substrate. (B) Tailoring the spectral response of the photodetectors fabricated on the SDC epilayer on the type Ib diamond substrates. The spectral is normalized by the incident photon flux.

253

2.7.5.5 Tailoring the photoresponse properties

10–6

10–23 270 nm

10–8

10–10

Responsivity (a.u.)

Photocurrent (A)

Hydrogen terminated

Ib substrate

10–12

10–24

Oxygen terminated

Ib diamond

10–26

10–28

Hydrogen terminated

Oxygen terminated 10–14 –10

(A)

–5 0 Applied voltage (V)

5

10

10–30 200 250 300 350 400 450 500 550 600 650

(B)

Wavelength (nm)

FIG. 2.7.5.7 (A) Photocurrent and (B) spectral response from type-Ib diamond substrate with O- and Htermination.

[33]. The oxidation of the H-diamond surface reduced the dark current to the noise level. By using the subsurface junction, one can treat the type Ib diamond substrate by using H-plasma to tailor the overall photoresponse. The DUV photocurrent can be enhanced by five orders of magnitude after H-plasma treatment of the Ib diamond substrate, as displayed in Fig.2.7.5.7A. A photocurrent gain as high as 100 or more was achieved. The dark current of the hydrogen-terminated diamond detectors can be controlled by controlling the H-plasma treatment time. Alvarez et al. developed a method to control the dark current through ozone treatment [21]. In addition, the spectral response was also greatly improved after hydrogen termination. The 210 nm/visible light injection ratio was improved to be as high as 104, as shown in Fig. 2.7.5.7B. The shape of the spectral response is quite similar to those of thin homoepitaxial diamond layers grown on the type Ib diamond substrates, with the appearance of the 270-nm shoulder. The two-dimensional holes on the surface of hydrogenated diamond are depleted by the nitrogen in the substrate, leading to a low dark current. Upon DUV light illumination, photogenerated electrons are trapped by the nitrogen, leaving excess holes in the surface layer. Such a space charge separation induces the photocurrent gain. The H-plasma treatment is quite simple, providing an alternative strategy to develop high-performance diamond DUV photodetectors. Three-dimensional photodetectors: Due to planar electrodes in MSM photodetectors, the electric field is inhomogeneous and substantially confined in the region close to the detector surface. A three-dimensional structure diamond photodetector was proposed to enhance the charge collection efficiency [34–40]. For instance, Lwakaji et al. fabricated a stacked-structure ultraviolet detector with high-quality undoped and B-doped homoepitaxial CVD diamond layers [34]. Under 210-nm illumination, the photocurrent increased with voltage increasing and it was about 100 times larger at 60 V than that at 5 V. The rise time was as fast as 1.2 ms. However, the fall time was lower than 12ms, originated to the deep level defects in the HPHT substrates. Alexander used a laser to introduce graphite column electrodes into the diamond to produce a three-dimensional diamond detector [36], as shown in

254

2. Semiconductor diamond

Beam orientation A Connected to current meter

Connected to bias voltage

AI Metallisation PCB

Graphitic channel

FIG. 2.7.5.8

Beam orientation B

Schematic of graphite column electrodes [36].

Fig. 2.7.5.8. The introduction of graphite column electrodes in single-crystal diamond could collect carriers generated in the bulk. These graphite column electrodes showed good ohmic contact characteristics, and the X-ray response current was up to 300 nA at 100 V bias voltage. Liu et al. fabricated a three-dimensional diamond ultraviolet photodetector via the down-top method [38]. A 30-nm thick tungsten (W) interdigitated electrode was patterned on the oxidized epitaxial diamond layer by standard lithography and magnetron sputtering technology. Then, the sample was put into the MPCVD chamber again for a second epitaxial growth. After the second growth, a 500 nm-thick SCD layer was grown on the diamond area. An ohmic contact was formed between tungsten and diamond after growth. According to the I-V characteristics curve, the photocurrent current of the device was 19 μA under 220-nm illumination, much larger than the dark current of 4.72 μA, revealing high sensitivity. According to the spectral response of device, the responsivity sharply increased from 350 nm and reached its maximum value at 220 nm, suggesting good selectivity for a threedimensional diamond UV photodetector. In the future, the dark current should be reduced.

2.7.5.6 Multicolor diamond photodetectors Multicolor optical sensing with high sensitivity at designed wavelengths can be applied in a variety of applications such as imaging, surveillance, optical communication, remote control, and target identification [41]. However, the typical diamond photodetector only detects the deep UV optical spectral band and cannot achieve the multicolor detection. For enlarging the response range of the diamond photodetector, Sang et al. combined diamond with other semiconductors in different bandgaps to fabricate a multicolor photodetector at first. Monoclinic gallium oxide (β-Ga2O3) with a bandgap of 4.2–4.9 eV also exhibits high intrinsic resistance and high spectrum selectivity. By integrating a β-Ga2O3 nanobelt on an intrinsic diamond epilayer, dual-wavelength solar blind photodetectors with a cut-off wavelength

2.7.5.6 Multicolor diamond photodetectors

255

FIG. 2.7.5.9 (A) Dark current and photocurrent dependence on the applied voltage characteristics of the Ga2O3/ diamond heterointegrated photodetector and diamond photodetector, respectively. (B) Spectral response of the two band photodetectors at 16 and 32 V. (C) Spectral response of the two-band photodetectors [41].

of around 280 nm were demonstrated. The heterointegrated β-Ga2O3/diamond device showed an extremely low dark current, and a high photocurrent-to-dark current ratio greater than 103 and 104 was achieved when the device was exposed to 220-nm (diamond band edge) and 240-nm (-Ga2O3 band edge) light illumination, respectively, as shown in Fig. 2.7.5.9A. High-performance dual-band photodetectors for visible light and DUV light detection can also be developed by integrating CdS (bandgap: 2.5 eV) nanowires on the diamond intrinsic layer. The spectral responses (Fig. 2.7.5.9B) of the CdS/diamond photodetector at bias voltages of 16 V and 32 V showed the two-band response to visible light and DUV light of CdS and diamond, respectively. A CdS (2.5 eV) nanowire and ZnO (3.37 eV) tetrapod-branched nanorod were integrated on diamond to fabricate three-band photodetectors covering light from the visible region to the UV-A region and the DUV region. The three-band absorption from the three materials can be distinguished in the photocurrent response spectrum (Fig. 2.7.5.9C), which indicated very good spectral selectivity. Liu et al. widened the spectral detecting range of the diamond photodetector via combining diamond with TiO2. A 450-nm-thick TiO2 film was directly deposited on the SCD epitaxial layer by the radio frequency magnetron technique. W electrodes were patterned on TiO2/diamond film to fabricate the UV photodetector. This device exhibited a 1.12-pA dark current at 30 V, and showed a two orders of magnitude UV-to-visible rejection ratio. Compared with that of the diamond photodetector, this device indicated increasing responsivity in a wide light wavelength range, which could be ascribed to the gradient energy band structure in the interface of the TiO2/diamond film. In addition, the device showed higher responsivity than that on diamond. Transient response showed that the increasing time of the device was 20 μs and the decreasing time was 1000 μs [42]. Because the thickness of the TiO2 used was about 450 nm, much deep UV light might be adsorbed. Thus, in further research, the same group also farbricated a TiO2/diamond photodetector by using an ultrathin TiO2 films with thicknesses of 18 nm for sample A, 12 nm for sample B, 6nm for sample C, and 2nm for sample D. The photodetector exhibited an obvious selectivity between the UV and visible regions. Two response peaks at 225 and 290 nm appeared in the spectral response, as shown in Fig. 2.7.5.10 for sample A, corresponding to the cutoff wavelengths of diamond and TiO2. The responsivity at 290 nm reduced with the thickness of TiO2 decreasing (Fig. 2.7.5.10B). This could be ascribed to the film thickness because ultrathin TiO2 film can only absorb less UV light. Moreover, Chang et al. enlarged the

256

2. Semiconductor diamond

FIG. 2.7.5.10 (A) Spectral response of 18-nm TiO2/diamond photodetector at 8 V. (B) Spectral response of Samples B, C, and D [43].

spectral response range of the diamond photodetector from deep UV to near UV through combining diamond with NiO film. The NiO/diamond photodetector showed good repeatability and a two orders of magnitude UV/visible rejection ratio. Also, the NiO/diamond photodetector has a higher responsivity and a wider response range in contrast to a diamond photodetector, which was ascribed to the energy band structure of the NiO/diamond film interface [44].

2.7.5.7 Photoconductive gain mechanism The achievement of photoconductivity gain requires the injection of carriers from the contact to the semiconductor. This is the case for the photoconductor with ohmic contacts. Normally, no photocurrent gain could be obtained for the photodiode at reverse biases. The gain for a photodiode can be achieved through an avalanche breakdown process or modification of the junction interface state [45, 46]. One can further improve the photocurrent gain by inducing deep defects, either in bulk or on the semiconductor surface. In such a case, a strong PPC often appeared simultaneously [47]. Photoconductive gain was found for both an SPD at reverse bias as mentioned above or an MSM photodiode fabricated on O-diamond with high boron concentration [48]. In either case, the interface barrier height might be reduced or the depletion width narrowed, resulting from the interface states. By using WC contacts on a nondoped thick (10 μm) diamond epilayer to form the MSM photodiode, a photoconductive gain was observed by changing the incident DUV light or applied bias magnitude. The following phenomena were found: (i) The gain was larger than unity at high biases and was accompanied by a slow time response, as shown in Fig. 2.7.5.11. (ii) the gain decreased when the measurement temperature was above 105oC at high biases (i.e., 32 V), and (iii) the photocurrent showed no dependence on the DUV intensity or the measurement

257

2.7.5.7 Photoconductive gain mechanism

(A) FIG. 2.7.5.11

(B) (A) Dependence of photoconductivity and (B) transient response on the applied bias [48].

temperatures at low biases (i.e., 5 V). Based on these facts, the gain was proposed to be related to charge trapping at the metal/diamond interface during DUV illumination. These charges led to the bending of the interface band and the formation of a thin interface barrier layer between the metal and diamond. This thin interface barrier layer thus brings forward hole tunneling from the metal to the diamond layer above a certain bias, which contributed to the gain. Thermionic-field emission (TFE) or field emission (FE) of majority carriers through the metal/diamond interface barrier was proposed as the reason for the photocurrent gain at high applied biases. The thermionic-field emission current through a thin interface barrier can be expressed as    q V + Vp JTFE ¼ Js exp ε0 where  1=2    A∗ T πqE00 q φB  V p exp  Js ¼ E0 k  1=2 φB V  Vp + 2 cosh ðE00 =kT Þ ε0 ¼

E00 E00 =kT  tanh ðE00 =kT Þ

E0 ¼ E00 coth ðE00 =kT Þ rffiffiffiffiffiffiffiffiffiffi qh Nt E00 ¼ 4π m∗ εs   kT Nv Vp ¼ log p q

(2.7.5.7)

(2.7.5.8) (2.7.5.9) (2.7.5.10) (2.7.5.11)

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2. Semiconductor diamond

FIG. 2.7.5.12

Fitting of photocurrent versus applied voltage curves by TFE and FE tunneling mechanisms [48].

Here, A* is the Richard constant (96 A/cm2 K2), q is the electron charge, h is Planck’s constant, k is the Boltzmann constant, m* is the effective hole mass, ΦB is the Schottky barrier height of the WC/p-diamond interface, V is the applied bias, εs is the dielectric constant of diamond, p is the hole density in the illuminated region, Vp is the Fermi-level potential in the diamond, and Nv is the effective density of states in the valence band. E00 reflects the tunneling probability. The deep trap with a concentration of Nt near the WC/diamond interface behaved as an acceptor. The field-emission tunneling current is expressed by a FowlerNordheim process as   b 2 +a (2.7.5.12) JFN ∝ V exp V where a and b are constants. These two injection mechanisms in Eqs. (2.7.5.7) and (2.7.5.12) were utilized to fit the I-V characteristics in Fig. 2.7.5.12, where ΦB ¼ 1.6 V, m*¼0.8m0, and εs ¼ 5.7 were used. For the JTFE current fitting, the hole density was 2.1 1017 cm3, and for the JFN current fitting, a and b were 23.9 and 61.4, respectively, for DUV intensity P¼ 15.6 μW/cm2, and 23.3 and 60.4, respectively, for P¼ 20 μW/cm2. It was found that thermionic field emission dominated the carrier transport at the metal/diamond interface when the DUV light power P was lower than around 5 μW/cm2. When P was larger than 15 μW/cm2, Fowler-Nordheim tunneling governed the photocurrent at the applied biases larger than 9 V. The fittings also revealed that the DUV illumination generated an additional charge of acceptor-type traps close to the interface, which was the origin of the thin interface barrier. For a photoconductor or a Schottky photodiode (forward biases) fabricated on borondoped diamond epilyers, a huge photoconductivity gain was observed [48]. Boron doping was a key factor to induce photocurrent gain and PPC upon the illumination of DUV light. An electron trap with a thermal energy of 1.37 eV was obtained, which was considered to be responsible for the gain and PPC. On the other hand, the optical quenching experiments showed a threshold at around 2 eV, which is related to the lattice relaxation of the trap.

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An absolute transient negative photocurrent (TNPC) was observed at zero or weak electric fields upon illuminating DUV light after optically quenching the PPC. The ionization of nitrogen in the substrate upon subbandgap light illumination speeded up the PPC quenching process and led to the absolute TNPC due to the transient electron filling of the charged nitrogen in the substrate.

2.7.5.8 Summary Diamond photodetectors are sensitive to DUV light and have a very low response to visible light. The performance of detectors based on PCD is influenced by grain boundaries, defects, and grain size. PPC is observed in PCD detectors due to the existence of minority carriers trapping the center in the grain boundary. Enhancement of performance of PCD detectors can be achieved through regulating device structure. Detectors based on SCDs show better performance and also have PPC caused by interface defects or a deep level in the SCDs. In order to develop high-performance diamond DUV photodetectors to satisfy the 5S requirement, one needs to obtain a high-quality diamond crystal or epilayer, engineer the impurities, control the interface, and properly design the device structures. When thick high-crystal quality and high-purity diamond is not available, one can tailor the epilayer structure and metal/ diamond interface to tailor the photoresponse properties of the diamond DUV detectors. By using a submicron-thick diamond epilayer on ta ype Ib diamond to form the subsurface junction, a high-performance diamond DUV photodetector was developed. Enhancement of collection efficiency for carriers might be realized through three-dimensional construction of detectors. To extend the applications of diamond detectors into more fields, the combination of diamond film with other semiconductors provides a promising route to tailor the spectral response.

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