Recent progress of boron nitrides

C H A P T E R

4 Recent progress of boron nitrides Xingwang Zhang, Junhua Meng Key Lab of Semiconductor Materials Science, Institute of Semiconductors, Chinese Academy of Sciences, Beijing, China College of Materials Science and Optoelectronic Technology, University of Chinese Academy of Sciences, Beijing, China

Boron nitride (BN) is a compound isoelectronic with carbon, and like carbon, can possess sp2- and sp3-bonded phases [1]. Structure schematics for the four primary crystalline BN phases, such as hexagonal BN (h-BN), rhombohedral BN (r-BN), cubic BN (c-BN), and wurtzite BN (w-BN), are given in Fig. 4.1. The most common phase of BN is h-BN, which is comprised of sp2 B-N bonds forming planar hexagonal networks that are stacked along the c-axis in an AA0 AA0 A configuration. The in-plane and c-axis lattice constants of h-BN are 0.250 nm and 0.666 nm, respectively, which are very close to the values of graphite (a¼ 0.241 nm and c ¼ 0.670 nm). The c-axis lattice constants for h-BN and graphite are equal to twice the interplanar separation due to the stacking. While graphite is a semimetal (zero bandgap semiconductor), h-BN is an insulator with a bandgap of 5.9 eV. Because of the weak interplanar bonding in both graphite and hBN, the planes can slide easily against each other, thus both materials are excellent lubricants. The weak interplanar bonding also accounts for the ability to insert atoms or molecules between the hexagonal layers, thereby forming novel and technologically valuable compounds. Hexagonal BN may be made by a chemical vapor deposition (CVD) process where boron trichloride, ammonia, and hydrogen are reacted together at high temperature. Another phase of BN with sp2 bonding is r-BN. Although it is usually referred to as a hexagonal structure, the primitive unit cell is a rhombus. The r-BN phase is physically similar to the carbon polytype called rhombohedral graphite. In this structure, boron and nitrogen atoms are also arranged in a planar hexagonal network, but the hexagons are stacked along the c-axis in an ABCABCA configuration. The lattice constants of r-BN are a ¼ 0.2498 nm and c¼ 0.9962 nm. Their values are comparable to those of rhombohedral graphite, namely a¼ 0.2456 nm and c ¼ 1.004 nm. Little is known about the physical properties of r-BN because it does not exist in crystals of adequate size and as a pure phase without the presence of h-BN for many experiments. The r-BN powers are usually produced from turbostratic h-BN under high temperature.

Ultra-wide Bandgap Semiconductor Materials https://doi.org/10.1016/B978-0-12-815468-7.00004-4

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4. Recent progress of boron nitrides

FIG. 4.1 Structure schematics for the four primary crystalline BN phases: h-BN, r-BN, c-BN, and w-BN.

The second most common phase of BN is the cubic zinc-blend structure, which consists of sp3-bonded boron and nitrogen atoms with {111} planes arranged in a three-layer ABCABC stacking sequence. Cubic BN has a crystal structure similar to diamond, except that while the diamond structure is comprised of two face-centered cubic (FCC) carbon cells, c-BN has one boron FCC cell and one nitrogen FCC cell. The lattice constant for c-BN is a ¼ 0.3615 nm, which is very close to that of diamond, a¼ 0.3567 nm. The bonding between atoms is entirely covalent in the case of diamond while c-BN exhibits ionicity between 0.26 and 0.48, depending on the ionicity scale used. Like diamond, c-BN is formed in equilibrium at high temperatures and high pressures (HTHP) from the hexagonal material. Diamond and c-BN have wide indirect bandgaps with energies of 5.5 and 6.4 eV, respectively. Because of the strong atomic bonding in c-BN and diamond, both materials are extremely hard and are excellent thermal conductors. These and other properties explain why c-BN is currently used for coatings on cutting tools, and is desired for wear- and corrosion-resistant coatings, heat sinks, and high-temperature electronic devices. Another phase of BN has the wurtzite structure (w-BN). The carbon analog is called lonsdaleite or hexagonal diamond. Wurtzite BN is comprised of sp3 bonds in which the atoms are arranged in a buckled hexagonal configuration and stacked along the c-axis in an AA0 AA0 A arrangement. The lattice constants for w-BN are a¼ 0.2553 nm and c ¼ 0.4228 nm, comparable to the lattice constants of lonsdaleite, a¼ 0.252 nm and c ¼ 0.412 nm. This material is very hard, but probably not as hard as c-BN. Wurtzite BN is formed at HTHP from turbostratic h-BN or r-BN, but seldom exists in pure form without a significant inclusion of sp2-bonding BN. Not much is known regarding the physical properties of w-BN because the single w-BN phase has not yet been produced. In addition, similar to graphitic carbon, sp2-bonded BN is often found in a disordered turbostratic form (i.e., t-BN). This is the form of sp2-bonded material most commonly observed in BN thin films. For t-BN, the two-dimensional (2D) in-plane order of the hexagonal basal

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4.1 Properties, synthesis, and characterization of c-BN

plane is largely reserved, but these planes are stacked in a random sequence and with random rotation around the c-axis. Turbostratic BN exhibits a broad and diffuse diffraction pattern that is distinct from that for h-BN and r-BN because of these disorders. Accompanying this decrease in stacking order is a slight increase in the (0002) interplanar spacing (0.35 0.36 nm).

4.1 Properties, synthesis, and characterization of c-BN 4.1.1 Properties of c-BN Cubic BN attracts widespread interest as a promising material for many potential applications because of its unique physical and chemical properties, such as the highest hardness lower than diamond, wide bandgap, high thermal conductivity, high electrical resistivity, chemical inertness, etc. [1]. The physical and chemical properties of c-BN are summarized in Table 4.1. For comparison, the corresponding properties of diamond are also listed in Table 4.1. TABLE 4.1

Physical and chemical properties of c-BN and diamond

Properties

c-BN

Diamond

Structure

Zinc blende

Diamond

Lattice constant (nm)

0.3615

0.3567

Cleavage

(011)

(111)

Bonding length (nm)

0.157

0.154

Density (g/cm3)

3.48

3.52

Bulk modulus (GPa)

367

435

Microhardness (GPa)

75–90

80–120

3500 (10.5 MPa)

4000 (13 MPa)

Expansion coefficient (10 /K)

4.7

3.1

Thermal conductivity (kW/mK)

1.3

2.0

Thermal stability (K)

1573–1673

833–973

Graphitize (K)

>1773

1673–2073

Chemical stability (with metal)

Better

Resistivity (Ωcm)

10 –10

1010–1016

Dopable type

p- and n-type

p-type

Dielectric constant

4.5

5.58

Refraction index

2.117 (583 nm)

2.417 (589.3 nm)

Bandgap (eV)

6.4

5.5

Melting point (K) 6

2

Bad 10

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4. Recent progress of boron nitrides

Because c-BN has the highest hardness next to diamond, it can be expected to show excellent wear resistance in a frictional environment, and thus it is a natural candidate for hard protective coatings. Contrary to diamond, c-BN is chemically inert in oxygen atmospheres (at temperature as high as 1573 K). In particular, c-BN can be used in environments where it comes into contact with iron-based metals because of its chemical inertness at high temperatures toward iron, cobalt, and nickel, for which diamond cannot be used due to its high chemical affinity with them. These properties make c-BN even more attractive for tooling applications in dealing with ferrous metals. In fact, the bulk form of c-BN is now widely used in manufacturing industries, for example, in cutting tools and grinding wheels of hard iron alloys. Because c-BN is transparent within the wider wavelength, another usage of c-BN will be in a protective coating for optical elements, particularly optical windows, for example, of ZnSe and ZnS. Cubic BN has a wide bandgap (Eg ¼ 6.4 eV), good thermal conductivity (13 W  cm1  K1), a low dielectric constant (7.1), and a high breakdown field (8 MV  cm1). Moreover, in comparison to diamond, which is only p-type dopable, c-BN is both n- and p-type dopable, which has been demonstrated at HTHP grown c-BN single crystals. Hence, c-BN has the more powerful potential for high-temperature and high-power electronic devices operating in harsh environments than that envisioned for diamond films. Additionally, it also is very suitable for UV light-emitting diodes (LEDs). In fact, it has been reported that the LED factured from doped c-BN crystals can function up to 950 K with a wavelength of 215 nm (5.8 eV) [2]. A 193-nm deep-ultraviolet (DUV) solar-blind photodetector based on high-quality c-BN films was also fabricated [3]. Besides, c-BN can be used as a high-quality insulating film and a heat sink for electronic devices because of high thermal conductivity, a thermal expansion coefficient close to that of GaAs and Si, a lower dielectric constant, a higher electric resistivity, and better stability. Development of field emitters is desired for vacuum microelectronic devices and fieldemission flat-panel displays. Due to their negative electron affinity (NEA), chemical inertness, and mechanical hardness, c-BN-based field emitters have received some attention as an electron source with a highly reliable performance.

4.1.2 Synthesis techniques of c-BN Film Cubic BN has been available since the 1960s as bulk material (small crystallites) formed by the HTHP synthesis technique established by Wentorf [4]. In this technique, mixtures containing boron and nitrogen are subjected to HTHP conditions in a metal capsule, with typical temperatures and pressures of 1800oC and 85,000 atmospheres (8.6 GPa). Two primary aims in the field are to reduce pressure and temperature for the process and find new catalysts to grow larger c-BN crystals. For instance, using Li3BN2 as solvent, the temperature and pressure for the synthesis of c-BN crystals can be reduced to 1500oC and 5.5 GPa, respectively. The fact that c-BN is a stable phase raises hopes that further decreases in pressure and temperature during synthesis are possible. However, the equipment of HTHP synthesis is considerably complicated, and the resulting crystallites are usually a few millimeters in size, which is too small for most industrial applications. The key to obtaining larger c-BN crystals of high quality is to control the supersaturation of BN in the solvent. Pure c-BN crystals

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FIG. 4.2 Single crystals of c-BN synthesized by the HTHP method by using (A) Li3BN2 solvent under 5.5 GPa and (B) Ba3B2N4 solvent under 5.5 GPa. Reproduced with permission from T. Taniguchi, S. Yamaoka, Spontaneous nucleation of cubic boron nitride single crystal by temperature gradient method under high pressure, J. Cryst. Growth 222 (2001) 549. Copyright 2001, Elsevier.

should be colorless; however, the color of c-BN crystals synthesized by the HTHP method has been reported to be between amber and yellow, indicating the presence of impurities. The color of the grown c-BN crystals was demonstrated to depend on the solvent: amber crystals were prepared by using Li3BN2 (Fig. 4.2A) while c-BN crystals obtained within a Ba3B2N4 catalyst system were colorless (Fig. 4.2B) [5]. Since the synthesis of c-BN thin films under low pressure was first reported by Sokolowski [6] in 1979, great efforts have been made to grow c-BN thin films. Since the 1990s, the successful growth of c-BN films has been reported using various ion-assisted physical vapor deposition (PVD) and plasma-assisted CVD methods, such as pulsed laser deposition (PLD) [7], radio frequency (RF) sputtering [8], ion beam-assisted deposition (IBAD) [9, 10], massselected ion beam deposition [11], molecular beam epitaxy (MBE) [12], direct-current (DC) jet-plasma CVD [13], radio-frequency plasma CVD (RF-CVD) [14], inductively coupled plasma CVD (ICP-CVD) [15], and electron-cyclotron resonance microwave plasma CVD (ECR-MPCVD) [16]. A simplified schematic illustration of the two basic processes, that is, ion-assisted deposition and plasma-assisted deposition, is shown in Fig. 4.3. Energetic-ion bombardment during film growth by irradiation with an ion beam or substrate biasing is essential and a common feature regardless of the methods of film synthesis. A review of c-BN preparation methods and their parameter spaces was presented previously [17–19]. 4.1.2.1 Ion-assisted pulsed laser deposition Initially, there have been several attempts to produce BN films by a simple unassisted PLD. However, it has only been unambiguously shown that h-BN or t-BN can be growth by PLD. For instance, by using a Q-switched Nd:YAG laser to ablate h-BN targets onto Si (100) substrates in an NH3 ambient gas, stoichiometric h-BN films were observed at substrate temperatures above 773 K. The synthesis of polycrystalline c-BN films using ion-assisted PLD was well established in 1990 by Mineta et al. [20]. They used a continuous CO2 laser to ablate h-BN targets while simultaneously bombarding the substrate with nitrogen ions from a broadbeam ion source, and produced mixed phase c-BN/h-BN films. Friedmann et al. [7] used 248-nm pulsed laser ablation with a Kaufman-type ion source run with Ar and N2 source

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4. Recent progress of boron nitrides

FIG. 4.3 Simplified schematic illustrations of the two basic processes for c-BN deposition: (A) ion-assisted deposition and (B) plasma-assisted deposition.

gases to produce films with a large fraction of c-BN. An ArF excimer laser (λ¼ 193 nm, τ ¼ 20 ns) was also used to ablate a sintered h-BN target, and the growing film was simultaneously bombarded by energetic nitrogen ions from the radio frequency (RF) plasma. Recently, Sajjad and Feng synthesized c-BN films by a CO2 pulsed laser deposition technique at a significantly low substrate temperature (450oC) using ferrous oxide nanoparticles as the catalyst [21]. In PLD, the film is formed by collecting the material that is ablated from a target with a high peak-power laser. Due to the energetic particles involved in the ablation process, PLD offers the possibility of depositing new phase material, such as c-BN, which differs in structure from the original target. In the intense plume generated during ablation, high-energy ions, radicals, and neutrals can be created, allowing for nonequilibrium chemical reactions both in the plume and on the substrate surface. There are many advantages of using PLD for c-BN film deposition. First, PLD can easily be combined with ion bombardment, reactive background gases, plasmas, and other techniques to optimize the deposition process. Furthermore, ion and deposition sources are separated in ion-assisted PLD, which permits independent control of the deposition rate, ion flux, etc. Drawbacks of PLD include limited potential for industrial scale-up and particle incorporation into the film. The particles, which originate from the target, can be minimized by using dense BN targets and mechanical choppers. 4.1.2.2 Sputtering Sputtering is commonly used in industry for depositing films. Early attempts to use RF magnetron sputtering to deposit c-BN thin films were largely unsuccessful. By enhancing ion bombardment during film growth, the films with high c-BN content can be deposited by sputtering. It was found that c-BN-containing films were synthesized only when a negative bias voltage was applied to the substrate. The use of magnetic fields to extend the sputtersource plasma out toward the sample can increase the ion flux on the film, and hence allow for

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353

a higher deposition rate. Both conventional magnetron sources combined with external magnets/coils and unbalanced magnetron sputter sources have been used to extend the plasma. The necessary ion bombardment can also be provided by a Kaufman-type ion source or with an ECR ion source coupled with a substrate bias. In general, RF magnetron sputtering has a lower deposition rate than DC sputtering. When insulating targets such as h-BN are used in sputtering, RF magnetron sputtering is required. Luthje et al. [22] have deposited films with high c-BN content from a B4C target with N2 added to the ambient by DC sputtering. Hahn and coworkers [8] have deposited films with high c-BN content from a hot boron target by DC sputtering because boron does become conductive at high temperatures (>1073 K). However, they found that the ion flux in a DC sputtering system was lower than that in RF sputtering. Also, Schutze et al. [23] were able to obtain films with a much higher c-BN content using RF versus DC sputtering, possibly also because of differences in the ion flux. Encouragingly, the c-BN films produced by ion-assisted sputtering are as good as any of the other techniques, and more importantly, sputtering is easily amenable to industrial scaleup. A primary drawback of sputtering is the inherent coupling of the ion flux and deposition flux, which makes it more difficult to independently vary, control, and measure these important parameters. 4.1.2.3 Ion beam-assisted deposition (IBAD) Ion beam-assisted deposition (IBAD) (or ion-assisted electron-beam evaporation) is one of the first techniques used in early efforts at c-BN synthesis. Since Inagawa et al. [24] synthesized nearly phase-pure c-BN films by IBAD, other workers [10] have obtained films with high c-BN content by this technique. In some studies, ions bombard the growing film with little or no bias applied directly to the substrate. The other processes involve striking a plasma discharge and extracting and accelerating ions to the biased substrate. The most configuration of IBAD employs ion beam sputtering of boron together with simultaneous bombardment by nitrogen and argon ions from a Kaufman ion source. The first source delivers Ar ions, which hit a pure boron target with an energy of 1 keV. The growing film is bombarded with a mixture of nitrogen and argon ions from the second source, hitting the film with a typical energy of hundreds of eV. Recently, Hirama et al. [12] reported the synthesis of c-BN films by ion beam-assisted molecular beam epitaxy (MBE), in which c-BN films were grown by electron-beam evaporation of boron with the irradiation of argon ions and atomic nitrogen radicals. The flux intensity of boron atoms was monitored and controlled by electron impact emission spectroscopy (EIES). Nitrogen ions were generated by an RF radical source, and argon ions were generated by an RF ion source with acceleration grids. The acceleration voltage of argon ions was varied in the range from 150 to 320 V, which determines the ion energy of argon ions. The argon flux intensity concomitantly changed with the acceleration voltage because of the design of the ion source. They found that the metastable c-BN phase can be grown on diamond as a result of the interplay between competitive phase formation and selective etching. It is particularly useful in that the boron deposition rate, bombarding ion energy, ion flux, and ion species can be controlled and measured independently during the IBAD process because the ratio of bombarding ions from the second source to the number of deposited film atoms is one of the key parameters for the growth of c-BN films. However, other processes such as condensation and thermal desorption, implantation of ions, recoil implantation of

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atoms deposited on the surface, and sputtering have to be considered in the IBAD process. Therefore, the c-BN growth regime is a complex function of many parameters. 4.1.2.4 Mass selected ion beam deposition (MSIBD) Cubic BN films could also be grown by the direct deposition of energetic boron and nitrogen ions using mass-separated ion beam deposition (MSIBD), as demonstrated by Hofs€ass et al. [11, 25]. They found that there are sharp threshold values of 125 eV for ion energy and 423 K for the substrate temperature to form the c-BN phase. The parameter windows of c-BN formation (125–1000 eV) for MSIBD are wider than those of other techniques. MSIBD can deposit films with high c-BN content with not only a lower threshold value of ion momentum per incident boron atom but also fewer ions per incident boron atom than the ion- or plasma-assisted deposition technique. In the MSIBD technique, the deposition parameters such as ion energy, ion flux ratio of different ion species, and substrate temperature are independently controllable. Furthermore, both boron and nitrogen species are deposited as singly charged energetic ions and no noble gas, other ions, neutral atoms, or molecules are involved. This relatively simple deposition process makes MSIBD the ideal tool to study the influence of the deposition parameters on c-BN growth, leading to a better understanding of the nucleation and growth mechanism. A nearly phase-pure c-BN film without inert gas ions was synthesized by this technique. 4.1.2.5 Plasma-enhanced CVD PECVD is one of the first techniques used to deposit predominantly c-BN thin films. In the case of PECVD, various deposition methods were employed, for example DC jet plasma CVD [13], RF-CVD [14], ICP-CVD [15], and ECR- MPCVD [16]. In most cases, ions from either a microwave (2.4 GHz) or an RF (13.56 MHz) plasma are extracted and accelerated to the biased substrate. A variety of source gases have been used to produce films with high c-BN content in literature. These source/carrier gas combinations are given as the following: B2H6 in N2 or NH3; BH3-NH3 in H2; NaBH4 in NH3; HBN(CH3)3 in N2; and B3H3N3(CH3)3 in N2. Unfortunately, diborane (B2H6), the most commonly used source gas in PECVD, is both toxic and explosive. Furthermore, CVD processes usually have higher substrate temperatures than PVD processes because they rely on thermally driven chemical reactions to crack the source gases, for example, substrate temperature ranges from hundreds to more than 1000 degrees for the PECVD process. Using less toxic and nonexplosive N-trimethylborazine (B3H3N3(CH3)3) as the source gas needs the substrate temperatures above 1020 K to produce c-BN films. In order to increase the quality of c-BN films deposited by the ICP-CVD technique, a timedependent biasing technique (TDBT) was developed. TDBT is a technique in which a relatively high minus substrate bias voltage is applied initially when B2H6 (dilute in He) gas is introduced into a deposition chamber. Then, the substrate bias voltage is continuously and gradually reduced to a final appropriate value for film deposition. Using this TDBT, BN films with more than 98% cubic phase can be deposited under optimized deposition parameters [15]. TDBT is also an effective approach for enhancing cubic-phase nucleation, even on a dielectric substrate [26]. For most CVD methods, the lack of an effective chemical reactant similar to hydrogen in CVD deposition of diamond films, that is, for selective etching of nondiamond phases and stabilizing the sp3-growing surface, results in the necessity of applying a high negative bias

4.1 Properties, synthesis, and characterization of c-BN

355

voltage during c-BN depositions. In the early 2000s, a marked progress in c-BN film deposition was made by the introduction of fluorine into the CVD process using a DC jet plasma CVD [13]. High-quality c-BN films with micrometer-sized grains and low compressive stress were successfully deposited at high rates [27, 28]. Due to a low bias voltage and a large precursor supply, a rather high deposition rate of 0.3 μm/min was obtained. Due to the lower energy required during the CVD process, the crystalline size of c-BN is usually larger and the crystalline quality much better than the c-BN films grown by ion-assisted PVD methods. However, the high gas pressure (6.7 kPa) and high substrate temperature (1000°C) conditions in DC jet plasma CVD may provide only limited applicability for electronic devices. CVD is a practical deposition technique that is amenable to industrial scale-up. Furthermore, CVD is generally inexpensive and can coat an irregular shape better than PVD techniques. However, the more variable contaminants could be involved in the films deposited by PECVD, and other drawbacks include a lower c-BN content than PVD techniques and the use of hazardous source gases such as diborane. In addition, in the plasma-based deposition process, the control and quantification of bombardment parameters are not as direct as in PVD.

4.1.3 Characterization techniques of c-BN Film Because c-BN thin film is typically very small grained and highly defective, the exact phase characterization of BN films requires a combination of several complementary techniques. Therefore, it is important to better understand various characterization techniques of BN films. In this section, various characterization techniques of BN films, including advantages and limitations, are discussed. 4.1.3.1 Phase identification by vibrational spectroscopy Being vibrational spectroscopies, infrared and Raman spectroscopies are sensitive to the bonding of BN films. Fourier transform infrared (FTIR) spectroscopy is a straightforward, nondestructive, and rapid analysis tool that is used widely for BN phase identification. The vibrational mode of bonding is governed by the crystal symmetry, and the selection rules determine the optical activity under a different source of excitations. For c-BN, the single, triply degenerate phonon is both IR and Raman active and is split into a transverse optical (TO) component at about 1065 cm1 and a longitudinal optical (LO) component at about 1310 cm1 by the ionicity of the crystal [29]. For near-normal incidence of c-BN thin films, IR reflection and transmission spectra only give a peak at 1065 cm1 corresponding to the TO component of c-BN, as shown in Fig. 4.4A. It was reported that w-BN has peaks in IR transmission at about 1085, 1125, and 1250 cm1, and the first peak is the most intense. Thus, the peaks at 1125 and 1250 cm1 can be used to distinguish w-BN from c-BN. For h-BN, the two infrared active modes have TO frequencies at about 1380 and 780 cm1 (Fig. 4.4A), which are due to the in-plane B-N bond stretch and the out-of-plane B-N-B bending mode of h-BN, respectively [30]. For r-BN, the two vibration modes are closely similar to the IR-active modes of h-BN. In addition, t-BN, which is poorly crystalline sp2-bonded, has the same two IR peaks as h-BN and r-BN, except broadened. So the IR spectra of h-BN, r-BN, and t-BN are quite similar, and it is difficult to distinguish them by FTIR. In summary, a reasonably symmetric IR

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4. Recent progress of boron nitrides

FIG. 4.4 (A) FTIR spectra of c-BN and h-BN films and (B) Raman spectra of bulk c-BN and h-BN crystallites.

peak around 1065 cm1 is the vibration fingerprint of c-BN thin films while peaks near 780 and 1380 cm1 are the fingerprints of sp2-bonded BN. One of the advantages of FTIR is the quantitative estimate of c-BN content. Several approaches to quantitatively estimate c-BN content have been presented by many researchers [7, 31, 32]. For both reflection and transmission analysis, it is common to assume that the volume fraction of c-BN is given by:  (4.1) qc ¼ I1065 ðI1065 + I1380 Þ where I1065 and I1380 are the normalized intensities of IR peaks at approximately 1065 and 1380 cm1, respectively. Friedmann et al. [7] have simulated the IR reflection spectra of randomly oriented thin films of c-BN, h-BN, and mixed-phase h-BN and c-BN on silicon substrates to estimate the relative IR sensitivity factor between the two phases. They found that the IR sensitivity factors are roughly equal for h-BN and c-BN films on silicon and the volume fraction of c-BN was approximately within that given by Eq. (4.1). For a metallic substrate [31], the intensity ratio of Eq. (4.1) needed to be multiplied by about 1.4 to estimate the cBN content. The principal deficiency of this approach is the use of the optical constants of highly crystalline c-BN and h-BN. Actually, though, the optical properties of c-BN and hBN thin films are significantly different from single crystal c-BN and h-BN. Jager et al. [32] showed that c-BN content may be overestimated based on transmission spectra by Eq. (4.1). They proposed a new approach for quantification based upon the IR transmission spectrum, that is, fc-BN/fh-BN ¼ 0.6 α(1065 cm1)/α(1380 cm1), where α is the absorption coefficient and fc-BN and fh-BN are the volume fractions of c-BN and t-BN, respectively. In fact, because the multiplicative constant 0.6 is nearly the density ratio of h-BN to c-BN, the absorption coefficients suggested by Jager et al. [32] approximately scale with mass density. The complex microstructure makes determination of the optical constants of c-BN and particularly h-BN problematic, and the precise quantification calculation of c-BN content is very difficult. In particular, quantitative analysis of the FTIR spectra of thick c-BN films (over 700 nm) is extremely tricky. It was reported widely that the FTIR peak of c-BN films is shifted to wave numbers above the single-crystal value of 1060 cm1 because of the compressive stress existing in c-BN films. Fahy [33] showed that the calculated peak shift is 34 cm1 per 1% isotropic strain,

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357

and for uniaxial strain, is 61 and 10 cm1 per 1% strain for phonon polarization parallel and perpendicular to the strain axis, respectively. In fact, McKenzie et al. [34] found that the IR peak of a c-BN-containing film shifted from 1081 to 1055 cm1 after stress was relieved when the film delaminated from the substrate. The position of the TO vibration mode ωTO is influenced by the relative lattice compression Δa/a0 which is caused by isotropic pressure p (in GPa): ωTO ¼ (1054.7 cm1) + (3.39)p. In the case of biaxial stress σ (it is true for thin films), the hydrostatic part p ¼ 2/3σ and the intrinsic stress in the c-BN film can be calculated from the shift of the c-BN peak, as the following: σ¼

ðωTO  1054:7Þ 5:09

(4.2)

Following Eq. (4.2), a stress-free c-BN film should produce a TO vibration mode in the IR spectrum at about 1055 cm1. When the film is under compressive stress, the peak position shifts to a higher wave number. Besides stress, it has also been shown that other many properties, such as deviations in film stoichiometry, contamination, film thickness, substrate properties, and crystal size, may strongly influence infrared absorption peak shifts [17]. Furthermore, the intensity ratio of 1380–780 cm1 depends on the orientation of the h-BN layers with respect to the substrate surface. Therefore, it is possible to also obtain information on the orientation of the h-BN basal planes from the IR data. IR analysis is a powerful but not absolute technique for BN phase identification. Silicon is used usually as a substrate to deposit c-BN thin films; however, SiOx has an intense IR absorption peak between 1050 and 1100 cm1, depending on the oxygen content. This SiOx absorption feature can be misinterpreted as the signature of c-BN. Fig. 4.4B shows the typical Raman spectra of a single-crystal c-BN and h-BN. Cubic BN clearly exhibits TO and LO vibration modes at 1056 and 1305 cm1, respectively, while only one peak at 1365 cm1 associated with the symmetric in-plane stretching (E2g) is observed for h-BN. Due to its inferior sensitivity, Raman analysis has not been widely used to study c-BN films. The two characteristic Raman signals of c-BN are very much dependent on the grain size and defect density. It has been demonstrated in c-BN films that the Raman peaks downshift to a lower wave number and become asymmetrically broadened substantially with increasing defect states and/or decreasing crystal sizes, whereas the compressive stress in the films induces a blue shift of the peaks [18]. Hence, it is very difficult to obtain a well-defined Raman scattering from current c-BN films, which have only small, highly defective crystallites. The appearance of c-BN Raman signals may serve as an indication for high-quality films. The full width at half maximum (FWHM) of the Raman peaks is also a measure of the crystallinity of c-BN films, and the smaller FWHM indicates the higher crystallinity of c-BN films. 4.1.3.2 Composition and stoichiometry analysis X-ray photoelectron spectroscopy (XPS) and Auger electron spectroscopy (AES) can provide composition and stoichiometry analysis of BN films. In the XPS spectrum of BN film, the B 1s peak is centered at 191 eV and the N 1s peak is located at 398 eV. From their integral areas (or the intensities ratio) and the sensitivity factors, the film stoichiometry can be calculated. In the energy above the main peak of the XPS spectra, the plasmon peak exhibits a significant difference between h-BN and c-BN. For h-BN, the bulk plasmon and the π plasmon loss peaks are present at 25 eV and 9 eV, respectively, away from their 1s peaks. However, for c-BN,

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there is only one satellite bulk plasmon loss feature at 25 eV away from the main peak, and no π plasmon peak for either the B 1s or N 1s, in contrast to what is seen in h-BN films. There are some disadvantages to the AES and XPS methods. First, the composition information of only the near surface of BN films can be obtained by AES and XPS, but the nearsurface composition of BN films is sometimes different from the bulk. Second, the use of ion sputtering to clear the surface or depth profile the sample can alter the composition of the films because boron and nitrogen may be sputtered at different rates. Finally, the electronbeam energies/currents commonly used in AES have been observed to damage BN, which can be overcome by using very low electron beam currents and employing special detection techniques. Rutherford backscattering spectrometry (RBS) is also useful to provide compositional information on c-BN films. Especially, it can obtain the compositional information through the entire film thickness. This technique is based on analysis of energy spectra of He+ ions or protons scattered in the direction opposite relative to the sample. It is quantitative without the need for reference samples, nondestructive, has a good depth resolution of the order of several nm, and a very good sensitivity for heavy elements. RBS measurements were done on a thick c-BN film in order to test whether Si atoms diffuse from the substrate into the BN films at high temperatures [35]. By applying RBS measurements, argon depth profiles were determined for differently prepared c-BN films, revealing inhomogeneous profiles with a preferential Ar incorporation into the h-BN layers close to the substrate. However, the RBS equipment is rather expensive, and complicated simulations are required to interpret data from BN films. Moreover, the accuracy for low-atomic-number BN films on particular substrates is not high enough because the intensity of the RBS signal is closely related to the mass of atoms. 4.1.3.3 Bonding character analysis Electron energy loss spectroscopy (EELS) in the transmission electron microscope is sensitive to bonding character, and has been used to identify BN phases [7]. The near-edge fine structure of the boron and nitrogen K ionization edges is a useful fingerprint of the bonding arrangement of BN films. The boron K edge in h-BN consists of a sharp peak centered at 191.5 eV, arising from the comparatively narrow energy bond consisting of π antibonding orbits in the sp2 hybridization of boron atomic orbits, followed by a broader peak at 198.5 eV from the 1s-σ* transition. However, in c-BN, the π* peak is absent and only a σ* peak at 197 eV is present, as shown in Fig. 4.5A. In the low energy loss regime, the energy of the plasmon peak exhibits a significant shift between h-BN (Ep ¼ 21.5 eV) and c-BN (Ep ¼ 30.1 eV), reflecting the large difference in density between the two phases. In addition, standard Auger spectrometers may easily be used to measure electron energy loss spectra in reflection geometry (REELS). In the case of BN films, the energy losses corresponding to the excitation of a volume plasmon have been reported to be clearly different for c-BN and h-BN, resulting in well distinguishable energy loss lines. As shown in Fig. 4.5B, the plasmon energy of h-BN is 25.5–26 eV and the value for c-BN is about 30 eV. Both near-edge X-ray absorption fine-structure spectroscopy (NEXAFS) and energy-loss near-edge spectroscopy (ELNES) can also provide useful information regarding the electronic bonding characters of BN films. For the sp2-bonded h-BN and r-BN phases, all spectral features for both B 1s and N 1s NEXAFS can be separated into two groups according to their

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FIG. 4.5 (A) Boron K-edge EELS spectra in the transmission electron microscope acquired from bulk h-BN and c-BN crystallites. (B) Differentiated REELS spectra of h-BN and c-BN film recorded with an Auger spectrometer.

angle dependence: features reflect transitions of a core electron into the states of π symmetry while all the remaining features correspond to transitions with a predominantly σ character. However, π states are absent in the pure cubic phase due to sp3 hybridization. Because NEXAFS is a synchrotron-based technique, its application will be limited compared to the widely available technique of FTIR spectroscopy. 4.1.3.4 Morphology and microstructure analysis Transmission electron microscopy (TEM), scanning electron microscopy (SEM), and atomic force microscopy (AFM) are useful techniques for providing detailed morphology and microstructural information about BN films. Because of the nanocrystalline dimensions of the ordered regions of both the sp2 and sp3-bonded phases in BN films, the conventional SEM and AFM are not suitable for microscopy and microstructure studies, but highresolution TEM (HRTEM) is a powerful technique for providing detailed morphology and microstructural information of BN films. Fringe spacing and angle can be measured directly from HRTEM images or in reciprocal space using Fourier transform methods. The c-BN phase can unequivocally be distinguished from the h-BN phase by determining the interplanar distances for the (111) c-BN (0.21 nm) and (0001) h-BN (0.33 nm) planes, respectively. The microstructures of BN thin films have been studied by means of cross-sectional HRTEM. The results showed that a thin amorphous layer was formed in the initial stage followed by a well-oriented t-BN interlayer, and finally the c-BN phase nucleated on the oriented t-BN layer. It is worth noting that HRTEM contrast is very sensitive to crystal orientation, beam tilt, and the objective lens focus and aberrations. So caution must be exercised in the interpretation of HRTEM images. In most cases, the use of an image-simulation program is a critical aid to the interpretation of HRTEM images. In the dark-field TEM technique, segments of the t-BN (0002) ring and the c-BN (111) ring are commonly used to form the image for the BN system. In such an image, background is the region of low electron intensity, while the regions of high electron intensity correspond to grains of the interest phase. Caution must be taken in the interpretation of such images because of the overlap of the c-BN (111) and the t-BN (10) rings.

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4.1.3.5 Structure analysis by diffraction methods Electron and X-ray diffraction (XRD) methods are usually used to analyze the structure of material. In the case of BN films, the weakly diffracted X-ray intensity from the thin, lowatomic-number BN films usually results in poor counting statistics. For most c-BN films with their grain size of only some nanometers, the method is additionally hampered by the large broadening of the Bragg peaks. Therefore, less significant information can be obtained by the standard XRD using a conventional θ–2θ geometry. To improve XRD counting statistics, the grazing incidence geometries are usually adopted. On the other hand, for the pure and thick c-BN films with high crystalline quality, the XRD θ–2θ spectrum as well as the rocking curves can be clearly displayed. Contrary to the XRD method, transmission electron diffraction (TED) has the advantage of significantly better counting statistics. However, the inelastic and multiple scattering processes and the relative imprecision in determining the camera constant complicate the analysis of diffraction intensities. For t-BN, the general {hkil} reflections are absent, and only reflections from the basal, {000l}, and prismatic, {hki0} (which is usually abbreviated as {hk}), planes remain. Actually, the most commonly observed peak is the basal plane {0002} reflection, and the spacing of the {0002} planes (d ¼ 0.36 nm) is typically larger in t-BN than in the highly crystalline h-BN due to the weak bonding between layers. For c-BN, the strongest peak is the {111} reflection (d ¼ 0.209 nm), and the c-BN grain size may be estimated from diffraction line broadening. However, many factors may complicate the characterization of BN films by diffraction methods. For the t-BN oriented with basal planes parallel to the substrate, its {0002} ring will be suppressed in the plan-view TED pattern, leaving only the 2D (10) and (11) rings. The remaining rings are close to the strong c-BN {111} and {220} rings, and could easily be misinterpreted as arising from c-BN. Furthermore, crystalline h-BN has several strong peaks that are close to c-BN. For example, the positions of the h-BN {1010} (d ¼ 0.217 nm) and {1011} (d ¼ 0.206 nm) reflections are close to that of the c-BN {111} reflection (d ¼ 0.209 nm). Similarly, the strong {0221} reflection (d ¼ 0.211 nm) from elemental rhombohedral boron is close to the c-BN {111} reflection.

4.2 Growth features and models of c-BN films In the last several years, c-BN thin films have been extensively investigated, and some significant progress has been achieved. A large variety of PVD and CVD methods have been applied in producing c-BN thin films. It should be noted that most successful methods for preparation of c-BN films are based on ion bombardment of the growing films. Therefore, in this section, the discussions concerning the growth mechanism of c-BN films will concentrate on only ion-assisted deposition.

4.2.1 Features and growth models of c-BN films via PVD All models of c-BN film growth should be based on the experimental data. Prior to investigating the models, some common features of c-BN deposition by PVD methods are summarized briefly in the following [17, 36].

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(i) Intensive ion bombardment (tens to hundreds of eV) is required for the formation of the c-BN phase. For medium ion energies, the synthesis of the c-BN film can be parameterized in terms of the ion momentum per depositing atom. At lower ion energies, other parameterizations may be more accurate. (ii) With increasing ion bombardment and substrate temperature, the growth rates decrease and the sequence h-BN ! c-BN ! nogrowth is found. In the case of c-BN, almost 40% of the incoming material is incorporated within the film. Low ion energies (Ei) can be compensated by high values of the ratio of ions to incoming boron atoms (F) and vice versa. (iii) There is a sharp low-temperature (about 423 K) threshold below which c-BN does not form. Tendentiously, high temperatures favor c-BN growth. Cubic BN deposition has been shown up to more than 1273 K. (iv) Nucleation of c-BN requires a higher ion energy or intensive ion bombardment than the subsequent growth of c-BN films. (v) Cubic BN films deposited on foreign substrates excluding diamond usually show a layered structure of substrate/thin amorphous BN (a-BN)/ preferentially oriented t-BN/c-BN. A continuous sp2-bonded a-BN/t-BN top surface layer is normally observed. (vi) Stoichiometric deposition is a basic requirement for c-BN growth. Films with high c-BN content are nearly stoichiometric. (vii) Cubic BN is generally nanocrystalline and consists of small, highly defective grains with typical crystal sizes between 20  100 nm. The c-BN crystallites have one set of (111) planes perpendicular to the substrate at the interface, but are otherwise randomly oriented. In addition to these experimental results, the phase diagram of the BN system has to be mentioned. Cubic BN is generally thought to be metastable with h-BN being the stable modification of BN based on the phase diagram proposed by Corrigan and Bundy [37]. However, recent experiments and calculations indicate that c-BN is the thermodynamically stable phase under ambient conditions. Regardless of the exact position of the thermodynamic equilibrium line, a significant kinetic barrier hinders the direct transition from sp2 to sp3 bonding under ambient conditions. Therefore, significant energetic-ion bombardment is favorable for the growth of the c-BN phase because c-BN is the denser modification and there is a kinetic or thermodynamic barrier for BN phase transformation. Based on these results, the opinion that ion-solid interaction is directly or indirectly responsible for the formation of c-BN has been widely accepted. Then, various growth models including quenching (thermal spike), selective sputtering, stress-induced formation, and subplantation have been proposed, as schematically shown in Fig. 4.6. However, the above models focus on a single aspect of the ion-solid interaction while they take into account the other effects only to the extent necessary to explain the experimental data. A completely satisfactory picture of c-BN formation has not been available so far. 4.2.1.1 Quenching (thermal spike) model The quenching model (also called the thermal spike model) developed by Weissmantel et al. [38] is the first description to explain the formation of c-BN. An ion doesn’t have sufficient energy to displace atoms near the end of its trajectory in a solid. The rest of the ion energy is transferred to the substrate by the creation of phonons in what has been termed a thermal spike. According to the theoretical analysis, the thermal spike can result in extremely high temperatures (up to several thousand K) and pressures (up to 10 GPa) during a very brief period of time. Weissmantel et al. [38] proposed that kinetics would be favored for the formation of the c-BN phase if the energy within the spike region is very rapidly quenched.

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FIG. 4.6 Schematic illustrations of (A) the quenching model, (B) the sputter model, (C) the stress model, and (D) the subplantation model.

In other words, the formation of c-BN proceeds via localized melting followed by rapid solidification. Spatial and temporal dimensions of thermal spikes can be estimated from the heat conduction equation. The deposition energy of 300 eV yields a time of 1011 s and a spike diameter of several nanometers; the latter is much smaller than the c-BN crystallite diameters of at least several tens of nanometers. Therefore, c-BN crystallites do not form from a single thermal spike. A simplified schematic representation of the quenching model is shown in Fig. 4.6A. Besides, a cylindrical thermal spike model was also proposed by Hofs€ass et al., in which ion deposition is treated as a cylindrical thermal spike with energy loss along the ion track, collision cascade effects, and conversion of energy into phonons and electronic excitations taken into account [39]. The model predicts that spike-induced atomic rearrangements appear to be crucial for the evolution of a cubic phase, but do not lead to density relaxation. The quenching model has not been formulated to an extent that is able to describe quantitatively the deposition parameters. Even a clear definition of the parameter space for c-BN deposition has not been performed. It is a major defect that the quenching model does not take into account the dynamic character of c-BN growth. If quenching without prior densification is sufficient, a transition from h-BN to c-BN should be possible by ion bombardment of h-BN targets. But many experiments indicate no effect on the BN phase by simple ion implantation into h-BN targets. The thermal spikes described above also play a role in both the static stress model and the subplantation models. In the quenching model, the thermal spikes are considered to be directly responsible for the phase transition. In contrast to the quenching model, however, in the stress and subplantation models, the thermal spikes are assumed to cause relaxation processes of structures produced by the collision cascade.

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4.2.1.2 Sputter model The sputter model developed by Reinke and coworkers [40, 41] assumes c-BN growth to proceed spontaneously by attachment of atoms to existing nuclei (crystallites) and the preferential sputtering of h-BN relative to c-BN. Because the sputter model assumes attachment to existing c-BN crystallites, it requires two preconditions: c-BN must nucleate by an independent process, and there must be an effective preferential bonding mechanism. The latter means that the incorporation of material is governed by the actual state of the growing film; the growth of h-BN on h-BN and c-BN on c-BN is favored. Simultaneously, the assumption of preferential bonding means that c-BN grows spontaneously, and the role of the ion bombardment is only to suppress h-BN growth. In the sputter model, the ion flux should be sufficiently high to etch all the deposited h-BN; only c-BN growth then occurs because of its lower sputter yield. A simplified schematic representation of the sputter model is shown in Fig. 4.6B. The sputter model can well describe the most observed experimental results. A major critical point is the surface nature of c-BN films. The sputter model suggests that sp3 bonding can only result if an incoming atom preferentially sp3 bonds on an existing sp3-bonded surface site. This contradicts the fact that the near-surface region of c-BN film is sp2 bonded. Some studies indicate that fewer and/or lower energy ions are required to grow c-BN once c-BN nucleates; however, selective sputtering cannot be the phase-selecting process for these conditions. Furthermore, the sputter model cannot account for the nucleation process due to the assumption of the preferential bonding mechanism. 4.2.1.3 Stress model High compressive biaxial stress of c-BN films in the GPa range is experimentally established. Several groups assume this stress σ to be responsible for the formation of the c-BN phase, and two different models (i.e., static and dynamic stress models) were developed. Fig. 4.6C shows a simplified schematic representation of the stress model. Static stress model: The static stress model developed by McKenzie and coworkers [34] suggests that c-BN forms because the ion-induced stress drives the BN material into the region of the c-BN phase diagram where the cubic phase is favored. McKenzie et al. [34] observed that the c-BN content increased sharply above a particular value of the compressive film stress (4  5 GPa), and this threshold stress was found to be insensitive to deposition conditions. Using the thermodynamic boundary between graphite and diamond, the equivalent boundary between h-BN and c-BN is estimated as 3 GPa, which is near the observed transition of c-BN formation (4  5 GPa). However, if c-BN and h-BN have the same ion-induced strain, c-BN will have a much higher stress because of its higher elastic modulus. The stress will necessarily increase because the effective modulus of the film increases with increasing c-BN content. Therefore, the stress in the film may be only a consequence of increasing c-BN content. Dynamic stress model: The static stress model considers static stress frozen in c-BN films after the thermal spike. In contrast, Mirkarimi et al. [42] considered that strain as well as stress are controlled by two factors, that is, dynamic concentration of interstitials and vacancies and the time-integrated accumulation of defects at sinks. These two factors control the instantaneous stress and the residual stress, respectively. In the dynamic stress model, ion bombardment creates a large instantaneous concentration of defects and thus a large instantaneous

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(dynamic) stress directly below the surface. Simultaneously, the temperature-induced relaxation processes still take place through diffusion and condensation mechanisms, which in turn determine the static stress level. Mirkarimi et al. [42] assumed that the maximum stress determines the formation of c-BN, and that a critical stress value is needed over some volume to induce the phase transformation. The static stress model contradicts most experimental parameter dependences while the dynamic stress model is in good agreement with the experimentally observed dependences. However, the critical point of the dynamic stress model is the temperature dependence. The substrate temperature Ts influences the defect concentration and stress, thus a sharp decrease of c-BN content should occur with increasing temperature. Furthermore, expansion of this model is limited by the lack of input parameters. 4.2.1.4 Subplantation model The term subplantation (i.e., subsurface implantation) was introduced by Lifshitz and coworkers [43] to describe low-energy ion implantation into near-surface regions. When ions with sufficient energy penetrate below the surface of sp2-bonded material, they would displace more sp2-bonded atoms than sp3-bonded atoms, resulting in an accumulation of sp3bonded sites. Fig. 4.6D shows a simplified schematic representation of the subplantation model. Their proposal relied on older studies indicating a significant difference in the displacement energies of sp2-bonded material and sp3-bonded material. However, it has been reported that differences in the respective displacement energies are too low to account for significant sp3 enrichment. Robertson [44] discussed the local densification by direct subplantation of ions. Sufficiently energetic ions penetrate the surface and enter interstitial positions, temporarily increasing the local density. This densification then causes the transition sp2 !sp3 because it is assumed that the state of hybridization adapts to the local density. Therefore, this model assumes a twostep deposition process: in the collision cascade, ions are rapidly condensed to form a dense, overcoordinated amorphous structure. The subsequent thermal spike causes a partial relaxation. Robertson’s subplantation model does not explain the observed energy and ion-flux dependence for c-BN formation. The subplantation model claims that the density increase directly results in sp3-bonded material. A more general perspective is that the ion irradiation raises the free energy of the sp2-bonded material until it is higher than the metastable sp3bonded material. Assuming a small kinetic barrier, the material then transforms. However, such free-energy arguments do not provide insight into the microscopic mechanism of the transformation. The subplantation model does not well describe the observed parameter dependence. For example, the subplantation model does not account for a sharp increase of c-BN content above some critical value of the ratio of ions to atoms. Perhaps a critical density is necessary to drive the transformation. Also, the subplantation model does not currently explain the sharp temperature threshold below which c-BN does not form. In summary, all these models concentrate on a single aspect of the processes of c-BN formation, and none of the simple models described above presents a completely satisfying interpretation of c-BN deposition. Particularly, both the quenching model and the static stress model explain only a small amount of the experimental observations. Also, the subplantation model does not well describe the observed parameter dependence. Only the sputter model

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and the dynamic stress model give a comparatively satisfying explanation for the observed experimental data. To explain all observed experimental data, a comprehensive model based on the sputter model [41] and the dynamic stress model [42] is proposed as follows. Initially, an amorphous SixByNz layer with a thickness of the ion range is formed due to the ion bombardment. Then, h-BN is deposited owing to a lack of c-BN nuclei. The h-BN layer is stable with respect to ion bombardment until the stress in the film reaches a certain value. The ion bombardment induces stress within this h-BN layer, resulting in a (0002) texture of the h-BN layer. This texture h-BN represents the thermodynamically most stable arrangement. By incorporation of material between the h-BN layers, the structure is densified. In this way, a transition from the h-BN (0002) to the c-BN (111) texture can take place without breaking existing bonds and without diffusion processes. Such a structure showing an almost perfect 2:3 lattice matching is favored by geometric similarities and has indeed been observed with TEM. Finally, these nuclei grow by attachment processes. Due to selective sputtering and the resulting lower growth velocity, the h-BN layer gradually becomes overgrown. The initial (111)-oriented growth of the remaining c-BN crystallites may be lost with increasing film thickness by secondary nucleation, for example, at grain boundaries where sp2 layers have been found experimentally. The size of the c-BN crystal is governed by the nucleation density. Therefore, the low crystal sizes reported experimentally are in agreement with a relatively smooth interface between h-BN and c-BN.

4.2.2 Growth features and models of c-BN films via CVD CVD refers to techniques that rely on surface reactions of gas-phase precursors to form the desired material. Because of the lack of an effective chemical reactant similar to hydrogen in the CVD of diamond films, applying a high negative bias voltage is required during cBN depositions, even for CVD methods. As a result, c-BN films deposited by such CVD methods show similar features as summarized for PVD films. Matsumoto and Zhang et al. [13, 16] showed the effectiveness of PECVD employing fluorine chemistry to deposit c-BN films. Fluorine works in many aspects as hydrogen in CVD diamond synthesis. Introducing fluorine into the gas phase dramatically reduces the substrate bias voltage needed for c-BN deposition. A key factor contributing to the effectiveness of PECVD methods is employing fluorine chemistry. In all those approaches, BF3 is introduced as a boron source, which decomposes into BFx clusters and fluorine atoms/ions. Some different experimental findings are observed during the growth of c-BN through the introduction of fluorine-containing species: (i) Fluorine can selectively etch the noncubic BN (t-BN and a-BN) phase, and the hydrogen-tofluorine ratio controls the deposition rate and phase purity. The selective etching by fluorine allows c-BN film deposition at very low ion energies. (ii) Unlike c-BN films produced by PVD, distinctive Raman spectra have been obtained from the samples prepared by fluorine chemistry due to the better crystallinity. (iii) The growth of c-BN is a surface process due to the absence of the a-BN/t-BN layer on top of the CVD-grown c-BN films. (iv) In the plasma, excited reactive gas species He, Ar, H, F, BFx, and NHx are generated. Among them, the BFx and NHx species are responsible for the growth of c-BN. (v) The c-BN film surface is borondeficient and the most dangling bonds are terminated by hydrogen atoms.

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Growth of c-BN using fluorine-assisted CVD is essentially different from the ion-assisted growth of c-BN previously described. On the basis of these conclusions, Zhang et al. [45] proposed the following sequence of c-BN growth from fluorine-containing species, as shown in Fig. 4.7: 1) The excited plasma contains reactive species, namely, H, F, BFx, and NHx, which are responsible for the etching of h-BN (mainly fluorine atoms) and c-BN crystal growth (BFx and NHx) (Fig. 4.7A). The c-BN surface contains a large number of nitrogen atoms terminated by hydrogen atoms, but a very small number of boron atoms terminated by fluorine atoms. 2) The c-BN crystal surface is activated by the impingement of plasma species and surfaceatom abstraction, thus providing free unsaturated sites for further growth (Fig. 4.7B). B-F bond breaking is most likely assisted by energetic specie bombardment (bias of about 20 V) while N-H bond breaking is easier. 3) Boron atoms are added to N-activated sites by incorporation of BFx species, whereas nitrogen atoms are added to B-activated sites by incorporation of NHx species (Fig. 4.7C). 4) Each boron surface site is subject to deposition by an NHx molecule at a much higher rate than the deposition of BFx on nitrogen surface sites (the N2/BF3 flow rate feeding the plasma is 100:3), thus leading to a B-deficient surface at any stage (Fig. 4.7D). It should be noted that the low kinetic energy ion bombardment is still needed for c-BN growth from fluorine chemistry because the ion bombardment is favored for the breaking of B–F bonds and the formation of a dangling bond needed for NHx adsorption. The relationship between the ion to boron flux ratio (Φion/B) and ion energy with and without the chemistry of fluorine is depicted schematically in Fig. 4.8 [46]. Without the chemistry of fluorine, the ion energy for c-BN formation must be above 40–50 eV and becomes higher with decreasing Φion/B. With the chemistry of fluorine, it is interesting to note that although the subplantation and momentum transfer models are unlikely to be applied, a similar tendency

FIG. 4.7 Schematic view of growth mechanisms of c-BN film in fluorine-based ECR-assisted MWCVD.

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FIG. 4.8 Schematic of the relationship between the ion to boron flux ratio (Φion/B) and ion energy with and without the chemistry of fluorine. Reproduced with permission from K. Teii, S. Matsumoto, Impact of low-energy ions on plasma deposition of cubic boron nitride, Thin Solid Films 576 (2015) 50. Copyright 2015, Elsevier.

is still seen, that is, the ion energy for the ICP and plasma jet is lower than that for microwave plasma because of the higher Φion/B. The ion energy for c-BN formation is below 40 eV, depending upon the plasma source. Two possible roles of low-energy ion impact (below some tens of eV) in the deposition are proposed by Teii et al. [46]. One is to enhance the migration of adsorbed species on the surface. The other is to produce available growth sites for incorporation of boron or nitrogencontaining species. In particular, low-energy hydrogen ions could abstract surface-bonded fluorine atoms for creating growth sites more efficiently than hydrogen atoms. A decrease in ion energy required for c-BN formation is thus preferable for improving the overall film quality. The possibility of growing c-BN via surface reactions without significant damage introduced by high-energy ion bombardment suggests that the growth of an electronic-grade cBN single is feasible.

4.3 Research progress on the growth of high-quality c-BN films In recent years, thin c-BN films have been deposited by a variety of experimental approaches. However, these c-BN films are of poor crystalline quality and composed of very small grains containing a high density of defects and grain boundaries that essentially prohibit electronic applications. On the other hand, the high compressive stresses developed during the deposition of c-BN films, induced by the intense ion bombardment, deteriorate the adhesive strength of c-BN films. Recently, two significant improvements in the growth of c-BN films have been achieved: low-stress thick c-BN film and epitaxy of c-BN on diamond, and both of them will be detailed in this section.

4.3.1 Stress relief and thick films of c-BN Energetic-ion bombardment during film growth by irradiation with an ion beam or substrate biasing has been found necessary to obtain the cubic phase, irrespective of the method

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of film synthesis. However, such ion bombardment is inevitably accompanied by the buildup of strong compressive stresses (5–20 GPa) until eventually, at a critical thickness, the c-BN films peel off. The maximum thickness of c-BN films is typically limited to a few hundred nanometers, and consequently, the application of c-BN coatings has been severely hindered. During the last decades, some postdeposition or sequential processes, such as annealing at temperatures above 800°C and ion irradiation, have been proposed to reduce the stress of c-BN films. More preferably, some attempts such as high-temperature deposition, simultaneous ion implantation, the addition of a third element, and a two-step deposition process have been undertaken to keep stress as low as possible during film growth. It has been widely reported that depositing a buffer layer between the substrate and the c-BN film is an effective process to reduce the compressive stress. In the 1990s, Okamoto et al. observed that the use of an intermediate boron layer and a graded BNx layer beneath the c-BN film prevented film delamination. Inagawa et al. reported that adherent c-BN films up to 1.5 μm could be deposited when the BNSi interlayer was used; nevertheless, no evidence was provided to show that the thick BN layers were mainly composed of the cubic phase. Yamamoto et al. [47] reported that adhesion of the c-BN films was vastly improved by a compositional gradient layer that contains boron, carbon, and nitrogen. In this work, a thin boron carbide layer (200 nm) was first deposited on silicon substrates. Following a boron carbide interlayer, a gradient layer with continuously increasing nitrogen content was prepared. After nucleation, the nitrogen content was increased to 100% and stable c-BN films grew to a thickness of more than 2 μm. However, the transfer of this technology to the other substrates required some further process modifications. Recently, the compressive stress of c-BN films with a compositionally graded B-C-N buffer layer was observed to rapidly decrease by the addition of hydrogen or methane into the Ar/N2 reactive gas during deposition [48]. Owing to the significantly decreased film stress, adherent micrometer-thick, cubic-phase dominant films can be allowed to form on silicon substrate. In addition to buffer layers, multilayer approaches have also been adopted to improve adhesion. For example, Noma et al. [49] designed a structure of the c-BN/t-BN/a-BN/B/Ti/TiN/Ti/substrate to prevent the delamination of the c-BN film from the substrate for the c-BN films deposited by a magnetically enhanced plasma ion plating (MEP-IP) method. It should be pointed out that the buffer-layer approaches may be suitable for mechanic applications, but not for the fabrication of electronic devices. A number of different approaches including postdeposition treatment and hightemperature deposition have also been undertaken to grow relatively thick c-BN films with good adherence to the substrate. It was found that films deposited at 950  1050oC took more than two months to delaminate whereas films deposited at temperatures <850oC delaminated in less than one week due to the reducing stresses. Fan et al. [50] reported that the stress in c-BN films can be reduced to a certain extent by using the two-stage deposition process with a c-BN seed layer grown first at 400°C followed by the deposition at elevated temperatures (800°C), whereas the c-BN content is nearly unaffected. Besides hightemperature deposition, postannealing can also reduce the stress in c-BN films. The compressive stress of the c-BN films reduces to 73% of the initial level, corresponding to a stress release of 2.2 GPa, by 1000°C postannealing during the first 20 min; it cannot be further relieved by prolonging the annealing time. The one drawback of these processes is the high deposition/annealing temperatures that will restrict potential applications to a few substrate materials.

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Boyen et al. [51] developed a way to prepare stress-relieved thick c-BN films by sequential ion-induced stress relaxation and growth. They found that the bombardment of high-quality c-BN films with 300 keV argon ions leads to a strong relaxation of their compressive stresses without destroying the cubic phase if the total ion fluence is kept below an upper limit. Furthermore, it was found that on top of such a stress-relieved film, a pure c-BN layer can be grown, but it builds up compressive stress again. Based on both results, a procedure was developed to grow 1.3-μm thick c-BN films exhibiting low residual stress and long-term stability under ambient conditions. Obviously, this method is a time-consuming and complex process that hinders its future practical applications. To simplify the sequential ion bombardment and growth, Fitz et al. [52] demonstrated that the intrinsic stress in c-BN films can be significantly relaxed during growth by simultaneous medium-energy ion implantation. The advantage of the present technique is that c-BN film can be deposited at optimum process parameters, which would usually cause high stress. However, because the boron evaporation rate became unstable after several hours of deposition, the c-BN films with a maximum thickness of several 100 nm were obtained in their experiments. As mentioned previously, more energetic conditions and a higher substrate temperature are generally required to initially nucleate than to subsequently grow c-BN. This two-step approach has been used in c-BN film synthesis for reduction of the radiation damage induced by ion bombardment during c-BN growth, improving the crystallinity, lowering the incorporated stress, and growing the thicker films. Accordingly, Litvinov and Clarke [53] have shown that the bias (nitrogen ion energies) necessary to initiate the formation of the c-BN on Si can be significantly decreased once a c-BN film of some minimum thickness (50 nm) is formed. The films grown with the reduced bias have less stress, and by adopting this two-step approach, they grew a 2.5-μm thick c-BN film at low bias (66 V) for 8.5 h. Actually, the lowest threshold energy of ion bombardment for c-BN film deposition can be reduced down to a few eV or less by using fluorine-assisted CVD, which will be discussed later. Another way of reducing the stress is the incorporation of a third element into the c-BN lattice. The idea is that adding the element offers additional possibilities to combine unstressed bonding angles and bonding lengths or to adjust lattice spacing. It has been reported that the incorporation of a small amount of carbon, oxygen, and silicon into the c-BN lattice leads to a reduction of the stress. Ulrich et al. [54] reported that the ultrahigh compressive stress was effectively reduced down to 3 GPa by controlling the incorporation of a small amount of oxygen into the deposition process. This consequently allows a 2-μm thick c-BN film to be grown on top of a coating system initiated with a boron-rich base layer followed by a nucleation layer with gradient chemical composition. Fan et al. [50] found that the incorporation of a small amount of Si (2.3 at.%) could result in a remarkable stress relief from 8.4 to 3.6 GPa; nevertheless, the incorporation of Si also led to a slight degradation of the c-BN crystallinity. Although several groups have made important progress in reducing the stress and allowing the deposition of c-BN films with a thickness above 1 μm, most of these approaches are rather complicated and not compatible with industrial needs. Significant progress was achieved by the fluorine-based CVD technique, resulting in a low-stress c-BN film with good adhesion with substrates. As discussed in Section 4.2.2, fluorine was recently found to be effective as an additional reaction precursor for c-BN thin film deposition, and it can selectively etch the noncubic BN phase [18]. The selective etching by fluorine allows c-BN film deposition

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at very low ion energies. Therefore, introducing fluorine into the gas phase dramatically reduces the substrate bias voltage (i.e., ion bombardment) required for c-BN deposition. Fluorinated CVD plasma induced by DC jet and low bias voltage enabled the preparation of 3-μm thick c-BN films, indicating relatively low internal stress. Similarly, introducing fluorine chemistry into ECR-MPCVD [16] or ICP-CVD [28] yielded thick and high-quality c-BN films. The film thickness was more than 10 μm for a 60-min deposition, as shown in the crosssectional SEM image in Fig. 4.9A [55]. Furthermore, distinctive Raman spectra, which cannot be observed for most nanosized c-BN films grown by ion-assisted PVD techniques, have been obtained from the samples prepared by the fluorine-based CVD technique, as shown in Fig. 4.9B [55]. The c-BN film deposited for 1 h shows a similar linewidth of Raman peaks to those acquired from the 4–8 μm c-BN HTHP single crystals, indicating that the crystallites in the film have comparable sizes.

4.3.2 Epitaxial growth of c-BN films As discussed in Section 4.2.1, c-BN is a promising semiconductor material for high-power electron devices operating at high temperatures because of its large bandgap energy of 6.4 eV, high breakdown field of 8 MV/cm, and chemical stability. However, up to now most of the cBN samples have been of poor crystalline quality and composed of very small grains containing a high density of defects and grain boundaries, making the prospect of using this material for high-temperature electronics an illusion. For electronic applications, single-crystalline c-BN thin films appear necessary, demanding an epitaxial growth on top of a suitable substrate. Since the 1990s, the heteroepitaxial growth of c-BN films has attracted

FIG. 4.9

(A) Cross-sectional SEM image of a 10-μm thick c-BN film deposited by DC jet plasma CVD. (B) Raman spectra of c-BN films deposited for 10 and 90 min as well as for 4–8 μm and 0.4 mm commercially available HTHP c-BN crystallites. Reproduced with permission from W.J. Zhang, S. Matsumoto, Investigations of crystallinity and residual stress of cubic boron nitride films by Raman spectroscopy, Phys. Rev. B 63 (2001) 073201. Copyright 2001, American Physical Society.

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great interest, and many groups worldwide have made a variety of attempts. Nucleating nanocrystalline c-BN thin films directly onto monocrystalline w-AlN (0001) films without a t-BN interlayer was first demonstrated in 2001 by Feldermann and coworkers [56]. Nanocrystalline BN grains with the cubic structure, and, more frequently, with the wurtzitic structure are found in direct contact with the AlN substrate. However, on many areas of the substrate, the well-known t-BN/c-BN layer sequence is present. Direct growth of c-BN onto the AlN substrate without the t-BN interlayer is observed on only one-third of the total substrate surface. The lattice constants of diamond and c-BN are 0.356 and 0.3615 nm, respectively, giving a weak lattice mismatch of 1.5%. Consequently, the best candidate as a substrate for c-BN epitaxial growth should be expected to be diamond. Actually, heteroepitaxial nucleation of c-BN on diamond single crystals and on CVD diamond has been demonstrated using the high-pressure high-temperature technique. Nevertheless, successful c-BN film nucleation onto diamond by low-pressure methods has very seldom been reported, although it has been considered in a number of studies. In 2003, there was a breakthrough in large-area heteroepitaxial growth of c-BN on top of diamond by IBAD [57, 58], resulting in c-BN layers with qualities suitable for electronic and optoelectronic device purposes. The cBN films were prepared on highly (001)-oriented diamond films of 10 μm thickness, which were grown on (001) Si substrates using MPCVD. The as-prepared diamond films exhibited a root-mean-square (RMS) roughness ranging between 40–60 nm (Fig. 4.10A). To reduce this value, the films were polished mechanically to a RMS roughness of 1 nm (Fig. 4.10B). The polished diamond substrate shows a very smooth surface with a few scratches. Fig. 4.10C shows the AFM image of a c-BN film grown on a polished diamond. It is clear that the grain morphology of the underlying CVD diamond substrate (Fig. 4.10A) is mirrored in the grain structure of the c-BN top layer, exactly implying the epitaxy of c-BN on diamond [57]. As shown in Fig. 4.10D, the FTIR absorption spectrum for a 30-nm thick c-BN film grown on a diamond substrate exhibits a unique narrow peak at 1075 cm1, indicating a 100% pure c-BN phase on diamond. Furthermore, the very narrow c-BN absorption peak with a FWHM of about 50 cm1, as compared with typical values of 150 cm1 in nanocrystalline samples, already implies that the c-BN film has a high crystalline quality. Ar depth profiles revealed by an RBS spectrum confirm the absence of an h-BN interlayer between the substrate and the c-BN layer. The HRTEM image (Fig. 4.10E) together with the in situ EELS spectra taken at the K-edge of boron and carbon (Fig. 4.10F) shows a sharp interface between the carbon and the BN layer and proves definitively that the apparently seamless crystallite is made of c-BN grown epitaxial on diamond [58]. Moreover, only c-BN (002) and (004) peaks can be clearly observed in the XRD spectrum, indicating that the c-BN (001) planes are oriented parallel to the diamond surface. According to the XRD results, the orientation relationship for the c-BN/ diamond heteroepitaxial system is determined to be c-BN (001)[100]jjdiamond (001)[100] [57]. The crystalline of epitaxial c-BN films is significantly improved by using single-crystalline (001) diamond as a substrate [58]. Epitaxially grown, 500-nm thick c-BN films are mechanically stable even under ambient conditions, though they still exhibit a compressive stress of 5 GPa, which opens exciting possibilities for electronic applications of c-BN films. In 2004, epitaxy of c-BN films on diamond was also accomplished by applying fluorine chemistry in an ECR MPCVD system [59]. The c-BN was grown on the as-deposited diamond films with no surface pretreatment, applying a substrate bias of 20 V. The c-BN films are

FIG. 4.10

AFM images for (A) an as-prepared CVD diamond (001) film, (B) a CVD diamond film after mechanical polishing, and (C) a 200-nm thick c-BN film grown on a polished CVD diamond (001) substrate. Scan areas in (A–C) are 1010 μm. (D) FTIR spectrum for a 30-nm thick c-BN film grown on a diamond substrate, demonstrating that only the cubic, and not the hexagonal phase (h-BN), is present. (E) Cross-sectional HRTEM image of the diamond/c-BN interface. (F) The two boron and carbon K-edge EELS spectra were taken in situ and serve to identify the diamond and c-BN phases. (A)–(D) Reproduced with permission from X.W. Zhang, H.-G. Boyen, N. Deyneka, P. Ziemann, F. Banhart, M. Schreck, Epitaxy of cubic boron nitride on (001)-oriented diamond, Nat. Mater. 2 (2003) 312. Copyright 2003, Nature Publishing Group. (E) and (F) Reproduced with permission X.W. Zhang, H.-G. Boyen, P. Ziemann, F. Banhart, Heteroepitaxial growth of cubic boron nitride films on single-crystalline (001) diamond substrates, Appl. Phys. A Mater. Sci. Process. 80 (2005) 735. Copyright 2005, Springer Nature.

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composed of small grains of several hundreds of nanometers with the underlying diamond grain facets still distinguishable. To understand the mechanism of epitaxial growth, c-BN films were grown on CVD diamond films at various substrate temperatures ranging from 420 to 1000oC by IBAD [60]. The results demonstrate that the epitaxial growth of c-BN films on diamond can be realized only at high temperature values (900oC) while at low temperatures, the same layered t-BN/c-BN structure is obtained, as commonly observed for c-BN growth on Si substrates. High substrate temperature is considered to be a key factor for the epitaxial growth of c-BN on diamond [60]. The presence of a t-BN interlayer at low temperatures is attributed to the amorphization or graphitization (i.e., radiation damage) of the diamond surface under the assisting ion bombardment during c-BN nucleation. When using diamond as the substrate material at elevated temperatures, the ion-induced damage can be sufficiently annealed out to maintain an ordered substrate surface during nucleation of the c-BN layer. Additionally, it should be noted that the resistance of the c-BN phase against destruction by ion bombardment is higher, probably due to the relatively high ionicity of the B-N bond. Therefore, the ion bombardment during c-BN deposition process allows to maintain the surface properties of c-BN. Recently, Hirama et al. [12, 61] achieved the heteroepitaxial growth of phase-pure c-BN (001) and (111) films on diamond (001) and (111) substrates by the ion beam-assisted MBE method with the irradiation of argon ions and atomic nitrogen radicals (N*). The ion irradiation for forming sp3-bonded BN is accompanied by ion etching during the epitaxial growth of c-BN films, although ion etching can be suppressed by lowering the acceleration voltage of argon ions. A V/III ratio higher than 1 and a growth temperature above 750°C as well as a suitable acceleration voltage are necessary for the epitaxial growth of c-BN films on diamond substrates. The microstructure of the heterointerface between diamond and c-BN has been investigated by using HRTEM. To determine the roles of the diamond interfacial layer on c-BN nucleation and growth, c-BN films were simultaneously synthesized on silicon and diamond-coated silicon in the ECR-MPCVD system. Some twinning of crystallites and small-angle grain boundaries was observed between the c-BN and diamond crystallites because of the slight lattice mismatch of 1.5%. The small-angle grain boundaries can be eliminated by imposing a little higher bias voltage during the initial growth stage. Chen et al. [62] applied the temperature gradient method to grow the c-BN single crystals on diamond at HTHP. By combining advanced HRTEM with first-principles calculations, they demonstrated that misfits at the interface between c-BN and diamond could actually be accommodated by a continuous stacking fault network that is connected by periodically arranged hexagonal dislocation loops. Furthermore, the energy band alignment of the c-BN/diamond interface was also determined by XPS measurements [63]. The c-BN valence band maximum (VBM) was 0.8 0.1 eV above the diamond VBM, which corresponded to the c-BN conduction band minimum (CBM) of 1.7 0.1 eV above the diamond CBM. Comparison with offsets predicted by theoretical calculations suggests that a C-N interface was obtained. Besides diamond, the direct growth of c-BN on commonly used Si (001) substrates has been attempted as well. By applying a hydrogen plasma pretreatment with proper substrate bias, c-BN was found to nucleate directly on Si (001) substrates in some local areas [64, 65]. Similar results have been reported by Zhang et al. [66], in which the elimination of an amorphous interface layer is due to the removal of the native Si-oxide layer by annealing at 1000oC under

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ultrahigh vacuum (UHV) conditions prior to c-BN growth. Thus, the epitaxial growth of c-BN on Si (001) substrates may also be possible in the future if appropriate pretreatment and deposition processes are applied.

4.4 Doping and electrical properties of c-BN Films Intrinsic c-BN exhibits a high resistivity and inconsistent conduction type depending on the preparation method. Because most electronic devices require controllable n- and p-type conductivities and low resistivity, much effort has been devoted to doping c-BN. In fact, both p- and n-type conduction in c-BN have been realized in bulk single crystals synthesized under HTHP, and a c-BN p-n junction has been fabricated by growing Si-doped c-BN on a Be-doped seed crystal [2]. However, the small size of c-BN crystals prepared by HTHP has inhibited its electronic applications. Therefore, doping of c-BN thin films and thus controlling their electrical properties are highly desirable. Probably due to the difficulty in growing high-quality c-BN thin films, comparatively few experimental results are available on the doping and electrical transport properties of c-BN films. Recently, thick c-BN films (>1 μm) with low stress were obtained by applying various methods [13, 16, 47, 51], and epitaxial growth of c-BN films on diamond has been realized by different groups [57–59]. Consequently, the doping of c-BN films has attracted research interest benefiting from progress in c-BN deposition. During the past decade, many different dopants were incorporated into c-BN thin films during the deposition process or by post ion implantation. Theoretical calculations and experimental studies indicate that substituted Be, Zn, or Mg at a B site results in p-type conductivity, whereas Si or S doping results in an n-type conduction [67]. In this section, theoretical predictions of band structure and doping in c-BN are first summarized. Next, experimental results on doping and electrical properties of c-BN films are critically reviewed. This is followed by a discussion of c-BN/metal contacts and p-n junctions.

4.4.1 Theoretical predictions of doping 4.4.1.1 Band structure of cubic boron nitride Theoretical approaches that have been used to calculate the electronic properties of BN can be generally divided into three categories: early semiempirical methods, ab initio methods based on density functional theory (DFT), and the latest post-DFT approaches such as hybrid-functional and many-body methods [1]. The overall nature and shape of bands obtained by the early theoretical studies are similar, but the quantitative differences are significant. One of the first ab initio calculations of c-BN was performed by Zunger and Freeman, producing an overestimated indirect bandgap of 8.7 eV. Development of the DFT and the practical approximations, such as the local density approximation (LDA) and generalized gradient approximation (GGA), have led to more consistent theoretical results. Several LDA-based theoretical calculations using different basis sets and pseudopotentials all yielded an indirect Γ-X bandgap with values ranging from 4.2 to 5.2 eV, leading to an underestimation of the bandgap. To reduce the unphysical self-interaction present in these approximations, Vogel et al. [68] introduced a modification to the LDA/GGA by using the

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self-interaction corrections method, obtaining an indirect bandgap of 6.1 eV for c-BN. Another alternative is a class of approaches based on hybrid density functionals, which produced improved bandgaps of 5.98 and 5.79 eV, depending on the basis sets. GW is one of the most widely applied theoretical approaches for electronic structure calculations. Earlier calculations for c-BN using a “one-shot” version of the GW method (G0W0) show an indirect Γ-X bandgap of 6.3 eV. Recent calculations [69, 70] using the one-shot G0W0 method resulted in bandgaps of 6.1 and 6.18 eV, which are in good agreement with the experimental results of 6.0–6.1 eV [71, 72]. Owning to neglect of the electron-hole attractive interactions, the bandgaps are typically overestimated by 10%–15% in GW approximation [70, 73]. Fig. 4.11 demonstrates the comparison of the band structure of c-BN calculated with LDA and self-consistent GW methods [69]. 4.4.1.2 Native defects [67] Intrinsic or unintentionally doped c-BN can exhibit n- or p-type semiconducting characteristics depending on the synthesis method, and the origin of conductivity is not yet clearly identified. To elucidate the possible mechanism of conductivity, the atomic geometries, the electronic structure, and the formation energies of native defects in c-BN have been calculated. The electronic structure of the B and N vacancies in c-BN has been investigated by several methods, including the full-potential linearized augmented-plane-wave (FPLAPW), the full-potential linear-muffin-tin-orbital (FP-LMTO), and the tight-binding linearized muffintin-orbital (TB-LMTO) approaches. All the methods give quantitatively consistent results. It is found that both boron and nitrogen defect states appear to form well-defined narrow bands in the forbidden gap of stoichiometric c-BN. The B-related vacancy band has a 0.4 eV width and is split off from the top of the valence band by 0.6 eV above the edge while the N vacancy can provide electrons thermally ionized to the conduction band and play the role of an effective donor impurity. These results suggest that B and N vacancies can be p- and n-type doping agents, respectively, which makes them the main candidates for the experimentally observed dopant compensation. Using first-principles total energy calculations, Orellana and Chacham [74] have also shown that N and B vacancies (VN, VB) are the most FIG. 4.11 Band structure of c-BN along the high symmetry directions. Thin green lines indicate the LDA calculation and thick red lines indicate the self-consistent GW calculation. A dotted line indicates the valence band maximum. Reproduced with permission from Ref. S.-P. Gao, Cubic, wurtzite, and 4H-BN band structures calculated using GW methods and maximally localized Wannier functions interpolation, Comput. Mater. Sci. 61 (2012) 266. Copyright 2012, Elsevier.

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stable defects in nonstoichiometric c-BN for p- and n-type conditions, respectively, and they also show intrinsic donor (VN) and acceptor (VB) characteristics. However, the VN in n-type c-BN shows high formation energies (>6 eV), which excludes it as a source of n-type conductivity. 4.4.1.3 Impurity doping In general, for the III–V compound c-BN, Group-II elements Be, Mg, and Zn substitute on boron sites as p-type dopants. Gubanov et al. [75] calculated the electronic structure and energies of Be and Mg in c-BN using the TB-LMTO method. They found that when Be and Mg substituted for B atoms, each created a delocalized level merged to the states at the valenceband edge. These partially occupied levels can result in p-type conductivity such as what has been observed experimentally. However, the characteristics of the hole states for the two impurities are rather different. Be-induced hole states appear to be strongly delocalized while Mg-induced impurity states are significantly more localized, and appear to be less effective for p-type doping than Be-induced states. Finally, the probability for Be and Mg impurities to substitute for N atoms in the c-BN lattice is calculated to be very small. The structural and electronic properties of Zn-doped c-BN were also investigated by Li et al. [76] via firstprinciple calculations based on DFT. The results suggest that Zn can substitute for both the B and N sites, both forming p-type conductivity of c-BN. Though Zn substitution of the B atom is energetically favorable, the large difference in electronegativity between Zn and N makes the acceptor levels strongly localized, thus reducing the effective carrier density. Zn substitution for the N atom of c-BN with a high p-type conductivity can only be achieved under conditions of rich B atom and N vacancy. Similarly, n-type doping in c-BN is possible by substituting group-VI elements (S or O) on N sites. Recently, the theoretical calculation for S-doped c-BN showed that S substituted for an N atom induces shallow donor levels near the bottom of the conduction band while S substituted for a B atom creates deep donor levels within the band-gap, both leading to n-type conductivity. Moreover, the total energy of S substitution at the B site (SB) is less than that of S substitution at the N site (SN), which means that SB is easier to obtain and more structurally stable. In contrast to S, an O impurity, either as an isolated substitutional defect or associated with a B vacancy, usually exhibits a negative effect on the conductivity characteristics of c-BN films [77]. Consequently, very few studies have attempted to dope c-BN with O. However, under B-rich conditions, the defect of O substituting at the N site (ON) can be either an acceptor or a donor dopant depending on the position of the Fermi level; thus, it can be considered as a source for either electron or hole compensation. Group-IV elements C and Si lie on the middle position between Group III and Group V in the periodic table, and thus are amphoteric impurities in c-BN; substitutions for both B and N must be considered. Early ab initio calculations using a supercell approach in connection with the FP-LMTO method predicted C-induced levels occurring at 0.5 eV and 5.1 eV above the valence band edge for C substitution at B (CB) and N sites (CN), respectively [78]. Orellana and Chacham [77] examined the energetics of CB and CN in c-BN by employing total-energy pseudopotential calculations, and they found that the substitutional impurity CB is a single donor while CN is a single acceptor. For a B-rich condition, both CN and ON have formation energies comparable to, or lower, than those calculated for vacancies, which are the lowestenergy intrinsic defects. The low formation energy and the acceptor character of CN suggest

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that this impurity might be related to observed p-type conductivities in unintentionally doped c-BN. Under N-rich conditions, the CB and CN defects have low formation energies for any Fermi-level position that is close in energy. Therefore, they would be equally favorable to be incorporated inducing simultaneous p- and n-type conductivities. Gubanov et al. [75] applied the TB-LMTO technique to investigate the electronic properties of isolated Si impurities replacing either B or N atoms in c-BN. They found that Si substituting at N sites (SiN) forms narrow, localized, and partially occupied levels within the bandgap, about 3.3 eV above the valence band edge and 1.0 eV below the conduction band edge. The large 3.3-eV energy separation between the partially filled impurity band and the top of the valence band makes electron transfer from the valence band impossible. Electrons occupying these states may provide the n-type activated behavior in Si-doped cBN. In contrast, a completely different situation is predicted for Si substituted in the B sublattice (SiB). The SiB impurity in c-BN induces highly delocalized states that overlap the conduction band edge, which contributes to n-type conductivity as observed experimentally in Si-doped c-BN. Most experimental studies on Group IV-doped c-BN have been based on Si doping while insufficient results have been reported regarding the efficacy of C as an n-type dopant. Besides Group II and Group VI, the electronic structures and properties of transition metal (TM)-doped and Group V-doped c-BN have also been investigated using DFT calculations [79]. The analysis of the band structure and density of states reveals that the B0.9375V0.0625N and B0.875V0.125N compounds exhibit half-metallic behavior, with a total magnetic moment of 2.0 and 4.0 μβ per cell, respectively. First-principles calculations with GGA revealed that the Al- and Ga-doped systems have the lower bonding energies of 11.544 and 5.302 eV, respectively. Difference charge density maps demonstrate that the electron loss increases after P doping and decreases after Al, Ga, and As doping.

4.4.2 Experimental studies of doping Bipolar carrier doping of both n- and p-types is a major obstacle for most wide bandgap semiconductors. For example, diamond can easily be doped as p-type with B, but its n-type doping is rather difficult. The situation is opposite for GaN and ZnO. Unlike other wide bandgap semiconductors, c-BN can be doped both n- and p-type, as demonstrated in bulk c-BN crystals [2]. Recently, impurities such as Be, Mg, Zn, S, C, and Si were incorporated into c-BN thin films during the deposition processes or by post ion implantation. Among of them, Be, Mg, and Zn have been identified as effective acceptors with an obvious increase in the conductivity of c-BN films [67]. On the other hand, to achieve n-type doping in c-BN films, some attempts have been devoted to the incorporation of S, C, and Si. In most experimental works, impurity incorporation is realized by in situ doping techniques via sputtering an additional target or evaporating with a standard effusion cell. Additionally, ion implantation is an alternative approach for doping semiconductor thin films and has a number of advantages, including precise control of dopant concentration, depth distribution, and doping area as well as high process reproducibility. Rapid thermal annealing is usually required after ion implantation to minimize irradiation-induced damage and to active implanted donors or acceptors.

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4.4.2.1 n-type doping S doping

The most promising dopants for n-type c-BN should be the Group-VI elements, and among these Group-VI elements, most studies have focused on the doping of c-BN with S. S-doped cBN single crystals were synthesized by the temperature-gradient HTHP method with a solvent of lithium boron nitride (Li3BN2) [80]. Hall measurements confirmed n-type conduction of S-doped c-BN crystals, and the temperature dependences of electrical resistivity clarified that the conduction is attributed to the thermal activation of carriers. The activation energy for conduction and the carrier concentration at room temperature are estimated to be 0.32 eV and 1014 cm3 for S-doped c-BN single crystals, and 0.47 eV and 1012 cm3 for unintentionally doped crystals, respectively, indicating an electrical doping effect of the S dopant [80]. However, a Hall mobility of 1 cm2/Vs at room temperature for the S-doped c-BN crystal is much lower than that of the unintentionally doped sample (20 cm2/Vs), but is comparable to values for Be-doped bulk p-type c-BN single crystals (2 cm2/V  s) [81]. In recent years, doping of c-BN thin films with S has been attempted by several groups. For example, S-doped c-BN films were prepared on p-type Si (001) and quartz substrates by introducing H2S into the ICP-CVD system, and the substrate temperatures were between 600 and 900 K. They found that nucleation of c-BN is suppressed with a very low H2S concentration while the growth of c-BN is less sensitive to S addition compared to nucleation. By predepositing a c-BN seed layer, S is successfully incorporated into c-BN films while maintaining a relatively high cubic phase content. In situ S-doping of c-BN films is also realized on Si (001) at 500oC using RF reactive sputtering by evaporating S into the working gas [82]. The carrier concentration in the S-doped BN films is approximately 6.5  1014 cm3. However, the cubic phase content is rather low in these films. For these in situ S-doped c-BN films, no definitive n-type behavior is observed by Hall measurements; only a rectification effect observed from an S-doped-BN/p-type-Si diode suggests the possibility of n-type doping of c-BN thin films. Moreover, the electrical conductivity of S-doped BN films is not reported in these studies. In addition to in situ doping, ion implantation has also been utilized to prepare S-doped c-BN films by two groups, resulting in highly conductive c-BN thin films. Zhang et al. [83] have attempted n-type doping of c-BN by S ion implantation, in which the c-BN film was deposited at 1173 K on p-type Si (001) substrates using IBAD. To produce a uniform depth profile of S ions in c-BN films, the implantation was carried out for multiple energies. It was previously reported that ion implantation was inevitably accompanied by radiation damage in c-BN, leading to a loss of cubic phase and a transformation to h-BN. To guarantee the preservation of the cubic phase throughout all implantation steps, the ion dose (5  1014 cm2) in this work is selected to be below the threshold for transforming c-BN into h-BN, which has been reported to be as high as 1015 cm2 for Ar and Si ion implantation. The slight degradation of c-BN crystallinity resulting from ion implantation can be recovered by postimplant thermal annealing. For a sample implanted with a total dose of 5 1014 cm2 and annealed at 1173 K, a dramatic reduction in electrical resistance by two orders of magnitude is observed while an Ar-implanted reference sample exhibited a resistance with the same magnitude as the as-deposited film [83]. These results indicate that the resistance decrease in an S-implanted c-BN film is primarily caused by S doping, not the defects resulting from

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FIG. 4.12 Electrical resistances (logarithmic scale) of the c-BN films after implantation and annealing versus the inverse temperature as expressed by 1/(kBT). The activation energy of the S dopant is estimated to be 0.28 eV from the slope of linear fitting. Reproduced with permission from X. W. Zhang, Z. G. Yin, F. T. Si, H. L. Gao, X. Liu, X. L. Zhang, Electrical properties of sulfur-implanted cubic boron nitride thin films, Chin. Sci. Bull. 59 (2014) 1280. Copyright 2014, Science China Press.

implantation-induced damage. Furthermore, the activation energy of the S dopant is estimated to be 0.28 eV from the temperature dependence of resistance, as shown in Fig. 4.12. The linear relationship between lnR and 1/(kBT) indicates that the conduction can be attributed to the thermal activation of carriers. However, it is still unclear whether S substitutes at the B or N site. In addition, the conduction type of the S-doped c-BN film cannot be determined by Hall measurements, owing to its high resistivity. Si doping

By introducing Si into c-BN, n-type conduction was first realized with HTHP-grown c-BN single crystals in the early 1980s; the activation energy of Si-doped c-BN crystals was determined to be 0.24 eV [2]. Recently, high-temperature thermal diffusion was also applied for doping Si impurities into HTHP-grown c-BN single crystals [84]. The activation energy of Si impurities was determined to be beyond 0.4 eV by the electric measurement. Due to the higher activation energy, it is impossible for Si impurities in c-BN to be completely ionized at room temperature. Therefore, the resistivity of Si-doped c-BN is still high, and the space charge limited current becomes the main conductive mechanism in c-BN. Because of the large radius of Si atoms and the small lattice constant of c-BN crystals, the diffusion depth of the Si impurity is restricted, and most Si impurities are distributed in the shallow layer underneath the surface of c-BN crystals. Besides c-BN single crystals, a number of groups have expended a great deal of effort in an attempt to realize n-type c-BN thin films using Si as a dopant. Similar to the doping of S, the formation of the cubic phase was easily disturbed by a high concentration of Si dopants. Ying et al. [85] prepared Si-doped c-BN films on Si (001) substrates at 400oC by IBAD via cosputtering of B and Si targets, and they found that the cubic phase content remained approximately constant for incorporated Si concentration up to 3 at.%. With increasing Si concentration, the c-BN phase was gradually disrupted and transformed to h-BN. When the Si concentration was increased to 6.7 at.%, the corresponding samples consisted primarily of h-BN with a small amount of c-BN. The difficulty in introducing Si atoms into c-BN lattice

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arises from the fact that the Si radius is much larger than that of B. It is essential to establish a doping method that does not destroy the cubic phase. By introducing Si at different stages (before nucleation, just after nucleation, and during the growth of pure c-BN phase), Oba et al. [86] found that the amount of Si dopant that suppresses c-BN nucleation is less than that which suppresses c-BN growth, which is similar to the case of S-doping. Furthermore, Yin and coworkers [87] found that the incorporation of Si into c-BN during film growth is impeded at high temperatures (900°C) by a strong tendency of Si to segregate toward the sample surface. This restricted the level of Si doping to approximately 200 ppm. However, Si segregation in the surface was not observed for Si-doped c-BN films deposited at 400oC in a similar work by Ying et al. [85]. A moderate film growth temperature is considered to be important to prevent Si segregation. As mentioned previously, Group-IV elements (Si and C) can substitute for both B and N in c-BN, and Si can substitute for B in c-BN as a donor dopant, leading to n-type conductivity. As in earlier works, Ronning and coworkers [62] found that Si is also bonded to B by replacing N, which may reduce its electronic doping efficiency. By ultraviolet photoelectron spectroscopy (UPS) and field emission electron energy distribution (FEED) measurements, no shift of the Fermi level toward the conduction band minimum was observed in the Si-doped c-BN films, and thus n-type doping can be excluded. This unsuccessful n-type doping may be related with the bonding configuration of Si in c-BN. In contrast, Ying et al. [85] reported that the incorporated Si atoms only replaced B atoms and combined with N atoms to form Si-N bonds. They found that the B/N atomic ratio decreased with increasing Si concentration while the (Si + B)/N ratio remained approximately unity for all Si-doped films, implying a substitutional doping of Si on the B site. A direct proof is provided by XPS spectra of Si 2p, B 1s, and N 1s core levels, as shown in Fig. 4.13. It can be seen from Fig. 4.13A and B that the binding energies remain at approximately 101.6 eV (Si-N bond) and 190.6 eV (B-N bond) for Si 2p and B 1s peaks for all the films, suggesting that the Si atoms combine with N atoms to form Si-N bonds in the films. However, the N 1s peak position shifts continuously toward lower energies with increasing Si concentration, as shown in Fig. 4.13C. Obviously, the N atoms combine with both Si and B atoms forming Si-N and B-N bonds, respectively, and the concentration of the Si-N bond increases with Si concentration. The local environment of Si in Si-doped c-BN samples was also investigated using X-ray absorption near-edge structure (XANES) and first-principles calculations [88]. The Si-K XANES of the Si-doped c-BN has a similar feature and almost the same chemical shift to a-Si3N4, implying that there are Si-N bonds and Si4+ in c-BN. According to the first-principles calculations, the substitutional Si at the B site is more favorable than that for the N site. Both the experimental and theoretical results confirmed that Si in c-BN prefers the B site to the N site. Bello et al. [89] grew Si-doped c-BN films by ECR-MPCVD using a gas mixture including SiF4. The resistivity of 250 Ωcm is obtained for the c-BN film with a Si/B ratio of 0.1, whereas the intrinsic c-BN resistivity is 1011 Ωcm. Hall measurements show an electron concentration of 2.1  1011 cm3 with a low Hall mobility of 1.18 cm2/V  s. Obviously, Si atoms replace B lattice sites to yield n-type conductivity. However, a detailed description of experimental results, such as cubic phase content and activation energy, is not presented. It has been reported that the conductivity of c-BN can be greatly improved by the addition of Si dopants. Ying et al. [90] have obtained Si-doped c-BN films by IBAD via cosputtering of B and Si targets at 400oC. They found that the resistance of Si-doped c-BN films gradually

4.4 Doping and electrical properties of c-BN Films

381

FIG. 4.13 XPS core level spectra, after surface cleaning, of c-BN films with Si concentrations from 0 to 15.7 at.%: (A) Si 2p, (B) B 1s, and (C) N 1s. The c-BN films were deposited on Si (001) substrates at 400°C by IBAD via cosputtering of boron and silicon targets. Reproduced with permission from J. Ying, X. W. Zhang, Y. M. Fan, H. R. Tan, Z. G. Yin, Effects of silicon incorporation on composition, structure and electric conductivity of cubic boron nitride thin films, Diam. Relat. Mater. 19 (2010) 1371. Copyright 2010, Elsevier.

decreases with increasing Si concentration. Rapid thermal annealing (RTA) conditions also have a significant effect on the electrical properties of Si-doped c-BN films, with the optimum RTA condition corresponding to a temperature of 1000oC for 3 min. The temperaturedependent electrical conductivities exhibit different slopes for the Si-doped c-BN films, suggesting that different conduction mechanisms dominate at different temperature ranges. The primary conduction mechanisms were extended-state conduction at lower temperatures, band tail-state conduction at intermediate temperatures, and short-range hopping conduction at higher temperatures. The resistivity of the 3.3 at.% Si-doped c-BN films decreases by only two orders of magnitude compared to undoped films. Thus, more effective methods are required to improve the doping efficiency of Si-doped c-BN films prepared by in situ doping techniques. The effects of the different deposition temperatures on the Si-doped c-BN epitaxial films have been systematically investigated by Yin and Ziemann [91]. They found that the Si concentration increases as the deposition temperature decreases, and the c-BN films doped at 420oC have lower resistivity than the ones doped at elevated temperatures. A negative Hall signal was observed, confirming n-type conduction for the Si-doped c-BN films. Furthermore, the temperature dependence of the electron transport properties indicates that the activation energy of Si-doped films is about 0.3 eV, and the increased Si concentration could improve the compensation from the deep-level acceptors. To minimize Si dopant segregation and enhance the electron transport in continuously in situ Si-doped c-BN films, multiple δ-dopings of single-crystal c-BN films heteroepitaxially grown on diamonds were performed [92]. For the purpose of δ-doping, the Si strip was pulsely inserted with the optimized sputtered area. Temperature-dependent carrier concentrations for a nominally undoped c-BN film, a continuously in situ Si-doped c-BN film, and a multiple Si δ-doped c-BN film are shown in Fig. 4.14A. It was found that, compared to the in situ Si-doped film, the resistivity of the multiple δ-doped film is reduced by two orders of magnitude at room

382

4. Recent progress of boron nitrides

900 750

Temperature (K) 450

300

nominally undoped in situ Si doped Si δ-doped

1017

Hall mobility (cm2/Vs)

Carrier concentration (cm−3)

1019

600

1015 1013 1011 109

(A)

15

20

25

30 1/kBT (1/eV)

35

continuously doped Si δ-doped nominally undoped

103

102

101 200

40

(B)

300

400

500

600

700

800

Temperature (K)

FIG. 4.14 Temperature-dependent (A) carrier concentration and (B) Hall mobility for a nominally undoped c-BN film, a continuously in situ Si doped c-BN film, and a multiple Si δ-doped c-BN film. The solid lines in (A) and (B) show fitting results. Reproduced with permission from H. Yin, P. Ziemann, Multiple delta doping of single crystal cubic boron nitride films heteroepitaxially grown on (001) diamonds, Appl. Phys. Lett. 104 (2014) 252111. Copyright 2014, American Institute of Physics.

temperature. The impurity band hopping conduction is considered to be the dominant conduction mechanism in the multiple δ-doped c-BN films. The Hall mobility for these c-BN films with different doping methods as a function of temperature is shown in Fig. 4.14B. A mobility of about 100 cm2/Vs at room temperature is obtained for the multiple δ-doped c-BN films, which is much higher than that of the normally continuously Si-doped c-BN film with a Hall mobility of 10 cm2/Vs. While most studies of Si-doped c-BN films have been based on in situ doping techniques, Yin et al. [93] recently reported a systematic study of Si-implanted c-BN thin films heteroepitaxially grown on diamond (001) substrates at 900oC. It was reported that the pure cubic phase can be conserved after Si implantation up to a concentration of 2  1020 cm3. n-type conductivity is confirmed by Hall measurements for these heavily Si-implanted c-BN films, and the sheet resistance decreases by seven orders of magnitude compared to undoped films. Due to the high Si doping levels, a semiconductor-to-metal transition is approached as revealed by the extremely small activation energies of typical 0.05 eV as well as by a nearly temperature-independent electron concentration up to 470 K. At high temperatures, an additional activated process sets in, resulting in a further increase in carrier concentration with an activation energy of 0.4 eV. 4.4.2.2 p-type doping Be doping

Based on first-principles calculations, Be is an ideal p-type dopant for c-BN. As early as 1987, Be was identified as an effective acceptor in bulk HTHP c-BN crystals with an activation energy of 0.23 eV [2]. According to the temperature dependence of resistance of Be-doped c-BN polycrystal, the activation energy of the c-BN polycrystals with Be concentration of 800 ppm and 1700–4500 ppm are determined to be 1.0 eV and 0.25–0.35 eV, respectively. Taniguchi et al. [94] also reported the varied activation energy of Be dopant in the range

4.4 Doping and electrical properties of c-BN Films

383

0.26–1.03 eV, depending on Be concentration. An increase in the Be concentration leads to a decrease in dopant activation energy in p-type c-BN. The Be concentration dependence of the activation energy is attributed either to that of an acceptor level or that of an interfacial level formed at grain boundaries. These results are in good agreement with the cathodeluminescence (CL) spectra of Be-doped HTHP c-BN single crystals, in which an increase in the impurity concentration leads to a shift in the position of the band maximum toward the short-wavelength range from 315 to 250 nm. Furthermore, it is proposed that several overlapping subbands, which are possibly associated with differently charged acceptor levels of beryllium, are located in the vicinity of the valence band of Be-doped c-BN. Besides HTHP c-BN single crystals, p-type conductivity of c-BN thin films was also achieved by Be ion implantation [95]. Prior to implantation, c-BN films were deposited at 950oC on polycrystalline diamond-coated Si(001) substrates by ECR-MP CVD. To yield a flat implantation profile, the implantation was conducted with several ion energies. For the sheet resistances for the sample implanted with a total dose of 2.5 1015 cm2 and annealed at 1420 K, a drastic reduction in resistivity by seven orders of magnitude is observed. As signalled by Hall measurements, the Be-implanted films exhibit p-type conductivity, corresponding to a hole concentration of about 6.1  1018 cm3 and a carrier mobility of 3 cm2/V s. It should be noted that the mobility obtained in this work is comparable to that of Be-doped HTHP c-BN bulk single crystals (2 cm2/V  s) [81]. The activation energy of Be dopants is found to be 0.20  0.02 eV, which is close to the value obtained in Be-doped bulk single crystals (0.23  0.02 eV) [2], but much smaller than that of Be-doped polycrystalline cBN (0.26–1.03 eV) [94]. These results suggest that Be implantation may be an effective p-type doping approach for c-BN thin films. However, it is worth noting that the use of Be metal presents technical challenges due to its toxicity. Mg doping

Magnesium is another dopant for achieving p-type c-BN. Lu et al. [96] reported the growth of mixed phase-BN thin films on Si (001) substrates using a neutralized nitrogen beam and electron beam evaporation of boron and its successful doping with Mg to achieve electrically active p-type BN thin films. It is found that all as-deposited BN films are p-type, and the Mgdoped BN thin films showed carrier concentrations in the range of 1.2 1018 to 5.2 1018 cm3 and mobilities of 18 to 52 cm2/V  s depending on the Mg evaporation cell temperatures. Both secondary ion mass spectroscopy (SIMS) and mass spectroscopy of recoiled ions (MRSI) analysis unequivocally confirm that Mg is incorporated in the BN films, and the activation energy for Mg in BN thin films is found to be 0.3 eV. However, the films are primarily h-BN with the c-BN content less than 30%, as revealed by FTIR spectra. More promising results were reported recently on Mg-doped c-BN films by Kojima et al. [97]. They deposited c-BN films on Si (111) substrates at 973 K by phase-regulated RF bias sputter, a thin Mg rod was inserted into the plasma, and this was negatively biased for in situ doping. The Mg concentration in BN thin films is controlled from 0.1 to several atomic percent by adjusting the position of the Mg rod. Less than 1 at.% Mg doping does not significantly affect the growth of the cubic phase; however, when the Mg content is higher than 1%, the incorporated Mg atoms lead to a transformation from the cubic phase to the hexagonal phase. It is found that the electrical conductivities of the c-BN films deposited in pure Ar atmosphere increase with increasing Mg concentration. A dramatic increase in conductivity by more than four orders of magnitude is obtained upon increasing the Mg concentration to

384

4. Recent progress of boron nitrides

2.1 at.%. The activation energies of Mg dopants are found to be rather low and decreased from 0.2 to 0.1 eV with increasing Mg concentration, indicating that shallow-level doping is possible in Mg-doped BN thin films. Films with 2.1 at.% Mg clearly reveal p-type conduction with a carrier concentration and mobility of 4  1014 cm3 and 6 cm2/V  s at 380 K, respectively. However, the fraction of nitrogen decreases with increasing Mg concentration up to 3 at.% for the films deposited in pure Ar while no modification of the composition and conductivity by incorporated Mg is observed when the c-BN films are prepared in Ar–N2 gas atmosphere. Thus, high conductivity by Mg doping is accompanied by a deviation from BN stoichiometry, indicating that the B/N ratio affects the electrical properties of Mg-doped c-BN films. Moreover, the Hall voltage is not reproduced reliably for c-BN films with low Mg concentrations. Zn doping

Besides Group-II elements Be and Mg, doping with Zn has also been attempted to achieve p-type c-BN. The Zn-doped c-BN films were prepared on fused silica substrates at 770 K by in situ RF sputter doping [98, 99]. The preparation process is similar to Mg doping of c-BN films as described above. For Group III–V semiconductors, Group II elements normally substitute on the lattice site of the Group III element. However, for c-BN films deposited in pure Ar atmospheres, the B/(B + N) ratio increases linearly with Zn concentration, implying that Zn preferentially replaces N atoms. This selective substitution can be attributed to the large difference in atom size between B and N [99]. Furthermore, lower dopant concentrations are achieved by Zn doping than by Mg doping using the same method, due to the difference in the stability of these elements in the c-BN phase. The conductivity of Zn-doped c-BN films deposited in pure Ar increased from 107 to 102 Ω1 cm1 at room temperature with increasing Zn concentration from 400 to 18,500 ppm, as shown in Fig. 4.15A. Nevertheless, it is

10

–1

600

500

Temperature (K) 400

300

0.6

Zn (ppm) 18500

10–2

18000

–4

10

9800 10–5 3100

–6

10

400

10–7 10–8 1.5

(A) FIG. 4.15

2.0

2.5

3.0 -1

1000/T (K )

3.5

Composition B/(B+N)

–1 –1 Conductivity (Ω cm )

0.58 10–3

0.56

A

0.54

B 0.52

0.5 0

(B)

1 0.5 Zn concentration (at.%)

1.5

(A) Conductivity of Zn-doped thin films plotted against the reciprocal temperature for the c-BN films deposited in pure argon. (B) Compositions of films of series A (deposited in pure argon) and B (deposited in a gas mixture of Ar:N2 = 18:2) measured by XPS as a function of Zn concentration. Reproduced with permission from K. Nose, T. Yoshida, Semiconducting properties of zinc-doped cubic boron nitride thin films, J. Appl. Phys. 102 (2007) 063711. Copyright 2007, American Institute of Physics.

4.4 Doping and electrical properties of c-BN Films

385

necessary to increase the Zn concentration to greater than 3000 ppm to achieve an obvious change in electrical conductivity with temperature. This can be ascribed to a high defect density in the films, which suppresses extrinsic electrical doping effects at low Zn concentrations. The n-type conduction is confirmed by Seebeck-effect measurements for high conductivity films; however, carrier mobility in these films has not been observed by Hall measurements. The Zn activation energy decreases from 0.3 to 0.1 eV with increasing Zn concentration, and these values are larger than those of Mg-doped films. Again, similarly to Mg doping, the selective substitution ZnN and the electrical doping effect are significantly suppressed by high N fluxes during preparation of c-BN films. As shown in Fig. 4.15B, doping is only effective for films with B/B+N 0.55 (deposited in pure argon), and the increase in conductivity is accompanied by a change in the B/N ratio. The primary experimental results on doping and electrical properties of c-BN thin films are summarized in Table 4.2.

4.4.3 Electrical properties and applications of c-BN films 4.4.3.1 Electrical contact [67] Degradation of device performance is often caused by high contact resistance between metal and semiconductor layers through thermal stress and contact failure. Formation of high-quality contacts is needed for electrical measurements and electronic application of c-BN. Therefore, investigations of how to fabricate appropriate electrical contacts with desired properties on c-BN films are of importance. In early studies, indium pellets were mechanically pressed onto the Mg-doped c-BN surface for the formation of electrical contact [96]. The samples were then annealed in an Ne atmosphere at 500°C for 30 min, and the annealing process was believed to be critical in obtaining adequate ohmic contact. In recent studies, metals such as Au, Ti, Ag, Al, and Cr were evaporated through contact masks in a vacuum deposition chamber yielding good ohmic contacts. For example, a proportional relation between the specific contact resistance and the specific resistance is observed for Ti contacts on Zn-doped p-type c-BN, indicating a typical metal/semiconductor contact [99]. To improve the quality of the contact, the samples are annealed at high temperatures from 600 to 773 K in vacuum or an N2 atmosphere. For most p-type c-BN films, Ag, Al, and Ti electrodes are used while Ti, Au, Cr/Au, Ti/Au, and Ti/Mo/Pt-Au contacts are deposited on n-type c-BN samples. Interface engineering, such as with a heavily doped surface layer, is frequently utilized for preparing an ohmic contact for wide bandgap semiconductors. However, for metal contacts on c-BN, only a few results have been reported on this topic. Yin et al. [100] investigated contact characteristics of Cr/Au layers, prepared by two different routes on c-BN thin films. Both pulsed laser ablation and thermal evaporation through a mask were used to fabricate the contacts. Linear I–V characteristics are obtained for all pad combinations indicating ohmic contact resistances. However, the evaporated Au/Cr contacts have to be ion bombarded additionally at room temperature with 300 keV argon ions to guarantee mechanical stability, whereas the Au/Cr films prepared by pulsed laser ablation are stable without further bombardment. The ion-induced interfacial intermixing at the metal/c-BN interface is believed to be the origin of the observed ohmic behavior.

386

TABLE 4.2 Experimental results on doping and electrical properties of c-BN thin films

Type Ntype

Dopant S

Doping method

Dopant concentration (at.%)

14

HPHT

6.510

Ion implantation Ptype

Be

2.110

Mg

Zn

11

1.18

0.32

Single crystal

[80]

Mixed BN film

[82]

0.28

c-BN film

[83]

0.24

Single crystal

[2]

c-BN film

[89]

1.17

c-BN film

[90]

0.35–0.40

c-BN film

[91, 92]

0.05

c-BN film

[93]

1–102

0.23

Single crystal

[2]

1

0.22

Single crystal

[81]

0.26–1.03

Polycrystalline

[94]

0.20

c-BN film

[95]

0.3

Mixed BN film

[96]

0.1

c-BN

[97]

0.1–0.3

c-BN film

[98, 99]

210

0.02 0.1

2

2.510 5

2.3

10 10 –10 13

16

100

20

10

17

HPHT

Ion implantation

Refs.

4

101–103

HPHT

HPHT

Crystal morphology

0.04

PECVD

Cosputtering

Activation energy (eV)

14

HPHT

Cosputtering

Resistivity (Ω cm)

510

2

10

0.29–0.08 0.1

Coevaporation Cosputtering

2.1

Cosputtering

2.0–0.04

18

610

3

41018

22

41014

6

7102

102–107

4. Recent progress of boron nitrides

Si

Mobility (cm2/Vs) 1

10

Reactive sputtering Ion implantation

Carrier concentration (cm23)

4.4 Doping and electrical properties of c-BN Films

387

Recently, deep-ultraviolet (DUV) solar-blind photodetectors were fabricated by depositing patterned Mo/Au (40/300 nm) electrodes on c-BN films [3]. The devices have a metal/semiconductor/metal configuration with circular interdigital electrodes. The I–V characteristics of the photodetector reveal a Schottky contact between the Mo/Au electrodes and c-BN films with a very low dark current. 4.4.3.2 Applications of c-BN Due to its excellent physical and chemical properties, c-BN is considered a promising wide bandgap semiconductor for high-power and high-temperature electronic devices operating in harsh environments. However, the device fabrication technology for c-BN material has not been explored extensively and it is limited primarily to p-n junctions fabricated from both intrinsic and doped c-BN. Unintentionally doped c-BN

Ronning et al. [101] investigated the conduction properties of unintentionally doped c-BN film deposited by MSIBD on p-type Si (001) substrates. They found that at low bias voltages, the I–V curves are symmetric with respect to the bias voltage and the Poole-Frenkel emission is considered to be the dominant conduction mechanism. The Poole-Frenkel emission describes electric field-enhanced thermal activation into the conduction band of electrons trapped in localized bound states. A dielectric constant of ε  8–10 and a trap potential Φ of about 1 eV are obtained from fitting I–V curves based on the Poole-Frenkel model. By applying high electric fields across the c-BN films, the conductivity rises step-like and the I–V curves change irreversibly from Frenkel-Poole to ohmic conduction. Up to now, there has been no satisfactory explanation for this irreversible behavior and the high bias part of the I–V curves. As mentioned previously, intrinsic or unintentionally doped c-BN can exhibit n- or p-type semiconducting character depending on the preparation methods. The I–V curves of p-n junctions between p-type Si (001) substrates and n-type c-BN films deposited by PECVD or RF sputtering exhibit typical diode characteristics. In other cases, p-type conductivity was observed in unintentionally doped c-BN films prepared by ICP-CVD and phase-regulated RF bias sputter. Rectification behavior is also exhibited in the I–V characteristics of p-type c-BN/n-Si heterojunctions, and a rectification ratio as high as 2.9  104 is achieved at room temperature. A thin t-BN initial layer between the c-BN films and substrates (except for diamond) is usually unavoidable and thus the initial t-BN layer may be critical to carrier transport through the c-BN film. Therefore, it is essential to elucidate the electrical transport properties of t-BN films and heterojunctions. Ronning et al. [101] reported that the electrical conduction of t-BN/steel structures could be well described by the Poole-Frenkel conduction mechanism. Rectification behavior is observed from the I–V characteristics of c-BN film with a thin t-BN initially fabricated by Nose et al. [102]. The rectification polarity is inverted in the double-layered film with thick t-BN, where conduction is found to be caused by Schottky and Frenkel-Poole emission conduction mechanisms, depending on the range of bias applied. However, different I–V characteristics of of p-type c-BN/n-Si heterojunctions were reported by Teii et al. [103] The t-BN/Si heterojunction shows no rectification while a typical diode characteristics is

388

4. Recent progress of boron nitrides

I V characteristics of (A) t-BN/n-Si and (B) c-BN/t-BN/n-Si heterojunction measured at different temperatures. The insets show schematics of the electrode configuration. Reproduced with permission from K. Teii, T. Hori, Y. Mizusako, S. Matsumoto, Origin of rectification in boron nitride heterojunctions to silicon, ACS Appl. Mater. Interfaces 5 (2013) 2535. Copyright 2013, American Chemical Society.

FIG. 4.16

achieved at room temperature for c-BN/Si with a thick ( 130 nm) t-BN interlayer, as shown in Fig. 4.16. In a further study, they found that a thicker t-BN interlayer increases the rectification ratio at room temperature up to the order of 105 at 10 V by decreasing the reverse leakage current [104, 105]. The higher bulk resistance of the thick t-BN interlayer impeded the transport of the minority carriers associated mainly with grain boundary defects in the films and, consequently, enhanced rectification. In addition to Si substrates, the I–V characteristics of t-BN/ZnO heterojunctions have also been reported [106]. Turbostratic BN/ZnO heterojunctions exhibit a pronounced rectifying behavior, low saturation current, and large ideality factors with n > 100. These unusual I–V characteristics are well described by a quantitative model using a serial connection of an ideal Schottky diode, a Poole-Frenkel type resistance, and an ohmic contact resistance. Doped c-BN

The successful fabrication of c-BN p-n junctions was first reported using HTHP single crystals. By growing an Si-doped n-type crystal epitaxially on a Be-doped p-type seed crystal at HTHP, Mishima et al. [2] fabricated a c-BN p-n junction diode that functioned from room temperature to 650oC. The luminescence spectra extended from 215 nm to the red, having a few peaks mainly in the UV region, and a similar emission was reported from a Be-doped p-type crystal that had localized p- and n-type regions in the same sample because of unintentional incorporation of oxygen impurities [80]. Besides c-BN homojunctions, Wang et al. [107] fabricated a heterojunction diode by epitaxial growth of B-doped p-type diamond film at 650oC on an Si-doped n-type c-BN bulk crystal using hot filament CVD. The I–V characteristics of the heterojunction diode indicate nearly perfect rectification behavior, and the turn-on voltage was determined to be as low as 0.85 V. However, the size of the c-BN crystals produced with the HTHP method was rather small, usually less than several mm.

4.4 Doping and electrical properties of c-BN Films

389

As discussed earlier, doping c-BN films has been attempted by many groups, resulting in both n- and p-type conductions. A rectification ratio of approximately 105 is obtained at room temperature for heterojunction diodes prepared by depositing S-doped c-BN films on p-type Si (001) substrates with the ICP-CVD method at temperatures between 600 and 900 K. Heterojunctions based on p-type c-BN films were also reported by He and coworkers [108]. They constructed a p-c-BN/n-Si heterojunction by implanting Be ions into c-BN films, which were deposited by RF sputtering on n-type Si(111) substrates at 600 K. The I–V curves of the heterojunction reveal a clear rectification feature, and the rectification ratio is determined to be about 200 at 8 V. The nonlinear surface I–V curves show the space-charge-limited current, indicating the existence of shallow traps in the Be-implanted c-BN thin films. Cubic BN is also a highly promising material for solar-blind deep-UV photodetectors capable of operating at high temperatures and in harsh environments. The fabrication of deepUV photodetectors based on c-BN thin films has already been reported [3, 109]. Soltani et al. [3] fabricated the solar-blind DUV metal-semiconductor-metal (MSM) photodetectors by depositing patterned Mo/Au electrodes on c-BN film, which is deposited on diamond substrates. As shown in Fig. 4.17, the DUV photodetectors present a peak responsivity at 180 nm with a very sharp cut-off wavelength at 193 nm and a visible rejection ratio (180 versus 250 nm) of more than 104. This indicates that c-BN-based detectors can detect DUV light directly without any visible and middle UV noise. In summary, wide bandgap c-BN has been considered a promising material for mechanical, electrical, and optical applications because of its unique properties. Although c-BN p-n homojunctions made from HTHP crystals were demonstrated decades ago, most of the early c-BN samples prepared as thin films have been nanocrystalline and exhibit high compressive stress, making the prospect of using this material for high-temperature electronics an illusion. Recently, substantial progress has been made in c-BN thin film deposition techniques, which opens up the opportunity for doping and electronic applications of c-BN thin films. Impurities such as Be, Mg, Zn, S, C, and Si have been incorporated into c-BN films, and the enhanced conductivity of c-BN films has also been observed by doping. However, despite much effort, electronic device fabrication based on c-BN is still lacking and only a few reports on c-BN p–n

FIG. 4.17 Spectral responsivity of a c-BN MSM photodiode measured at different bias voltages under the irradiation of a DUV monochromatic source in the range of 175–250 nm. Reproduced with permission from A. Soltani, H. A. Barkad, M. Mattalah, B. Benbakhti, J.-C. De Jaeger, Y. M. Chong, Y. S. Zou, W. J. Zhang, S. T. Lee, A. BenMoussa, B. Giordanengo, J.-F. Hochedez, 193 nm deep-ultraviolet solar-blind cubic boron nitride based photodetectors, Appl. Phys. Lett. 92 (2008) 053501. Copyright 2008, American Institute of Physics.

390

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junctions are available. In most studies, these heterojunction diodes are fabricated on Si (001) substrates and exhibit good rectification behavior. Using substrates with close lattice parameters is essential for depositing high-quality c-BN thin films because the lattice mismatch causes extended defects with detrimental effects at the interface. Recently, epitaxial growth of pure c-BN films with low stresses has been achieved on diamond (001) substrates. Similar structure and physical properties enhance the prospect of creating successful heterojunction diodes based on c-BN and diamond films, which would have very interesting electrooptical properties because of the large energy bandgap and an unusual type-II band alignment. Thus, more attention should be paid to the development of c-BN/diamond heterojunctions in the future. It is anticipated that continuing achievements in deposition and doping technologies will soon enable the realization of c-BN-based electronic devices.

4.5 Properties, synthesis, and applications of h-BN Film Hexagonal BN crystallizes similar to graphite in a hexagonal sheet layered structure, and therefore it is often referred to as “white graphite.” Within each layer of h-BN, boron and nitrogen atoms are bound by strong covalent bonds, whereas the layers are held together by weak van der Waals forces. The different types of bonding in two directions lead to a high degree of anisotropy for most properties of h-BN. Although the in-plane B-N bonds are dominantly covalent in nature, they show partial ionic behavior because of the difference in electronegativity between B and N (2.04 for B and 3.04 for N). As a result, h-BN shows less electrical conductivity compared to graphite and its increased interlayer interaction results in higher mechanical hardness. Due to the higher electronegativity of nitrogen, the π-electron is located at the nitrogen and therefore h-BN is an electrical insulator with a large bandgap. Furthermore, h-BN is stable in air up to 1000oC, under vacuum up to 1400oC, and in an inert atmosphere it can be used up to 2800oC. Hexagonal BN films have a wide range of attractive properties, including high temperature stability, a low dielectric constant, high mechanical strength, a large thermal conductivity, high hardness, and high corrosion resistance, leading to a number of potential applications as both structural and electronic materials. First, h-BN powder is traditionally used as a solid lubricant with low friction in numerous applications. It can be also added to a liquid to get dispersions with lubricating properties. Compared with graphite, the hBN can be used as a lubricant in an oxidizing atmosphere up to 900oC as well as at extremely low temperatures. Second, the excellent insulating and dielectric properties of h-BN combined with the high thermal conductivity make this material suitable for a large variety of applications in the electronic industry. As an insulator, it has a dielectric constant of 4 and a dielectric strength almost four times higher than that of alumina. As a thermal conductor, h-BN exceeds almost all other electrical insulators while maintaining high strength and low thermal expansion. Therefore, h-BN has also been applied as a charge leakage barrier layer for electronic devices and equipment. Moreover, h-BN is also used as windows in microwave apparatus and as insulator layers for metal-insulator-semiconductor field-effecttransistor (MISFET) semiconductors. Third, because of the nonreactivity with the melt and the nonwettability, pyrolytic h-BN is widely used for crucibles as well as linings and covers

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of heating elements in the equipment for semiconductor single-crystal production (e.g., Czochralski apparatus, Bridgman apparatus). Besides, h-BN coatings are used in the steel and iron industry to enhance corrosion resistance as well as to reduce the wear on sliding parts such as the crankshafts for compressors. On h-BN coated preforms, metal layers can be deposited, allowing one to get the free-standing metal sheets by easily peeling off. Fourth, lasing action in the deep UV region (225 nm) by electron beam excitation has been demonstrated in small h-BN bulk crystals synthesized by an HTHP technique, raising its promise as a deep UV photonic material. Recently, deep UV photodetectors based on h-BN films have also been fabricated. Finally, because 10B is one of the constituent elements in h-BN with a large thermal neutron absorption cross-section, solid-state neutron detector (SSND) fabrication using h-BN is another application of interest.

4.5.1 Band structure of h-BN Most theoretical calculations based on the first principles predicted that the bandgap properties of h-BN will show indirect transition properties. The methods based on DFT have predicted that the CBM is located at the M-point of the Brillouin zone while the VBMs are predicted at various points. For instance, LDA calculations, which usually produce underestimated values of bandgap by 1 eV, show h-BN to be an indirect bandgap semiconductor with a bandgap of 3.9 eV between VBM at the H-point and CBM at the M-point. By using a more accurate quasiparticle G0W0 method, the bandgap is increased to 5.4 eV, and because self-energy is weakly k-dependent, the band structure (indirect K–M bandgap) is similar to that of LDA in other aspects. Recent calculations for h-BN using a G0W0 approach show an indirect K–M bandgap of 5.95 eV, which is in good agreement with the experimental value of 5.955 eV. Fig. 4.18 shows typical quasiparticle band structure of h-BN compared to the band structure obtained with LDA [110]. Besides h-BN with AB stacking, several polytypes of h-BN are also calculated using dispersion-corrected DFT and G0W0 methods,

FIG. 4.18 Band structure of bulk h-BN along the high-symmetry directions. Solid lines and open circles indicate results of the LDA and GW calculations, respectively. The energy scale is relative to the top of VBM located at the T1 point. Reproduced with permission B. Arnaud, S. Lebegue, P. Rabiller, M. Alouani, Huge excitonic effects in layered hexagonal boron nitride, Phys. Rev. Lett. 96 (2006) 026402. Copyright 2006, American Physical Society.

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resulting in similar total energies of different hexagonal structures and indirect gaps of  5– 6 eV, only slightly smaller (50–150 meV) than the direct gaps at K-point. In contrast to the theoretical calculations, the experimental data on the bandgap nature of h-BN are contradictory; both direct [111] and indirect [112] bandgap have been reported, with the bandgap energies ranging from 3.6 to 7.1 eV. In general, the properties of the band edge can be determined by measuring the exciton-related luminescence and absorption spectra near the edge. However, neither band-edge exciton luminescence nor the absorption spectrum had so far been measurable because of the poor crystalline quality of h-BN samples. To determine the intrinsic bandgap nature of h-BN, the purity of the samples is particularly important because the optical properties near the bandgap can be sensitively modified by the impurity states. In 2004, the exciton-related luminescence at 215 nm was first observed by Watanabe and colleagues for single crystalline h-BN synthesized by the HTHP method, supporting the direct nature of the bandgap [111]. The intense band at 5.765 eV (215.0 nm) is the free exciton luminescence showing a Stokes shift from the position of the 1s absorption peak at 5.822 eV. A series of s-like exciton absorption bands from 1s to 4s appeared together with the plateau of the Elliott step, allowing estimates for the bandgap and exciton binding energies of 5.971 eV and 149 meV, respectively. In 2007, Kubota et al. [113] synthesized high-purity h-BN crystals at atmospheric pressure by using a nickel molybdenum solvent, and the obtained h-BN crystals emitted intense 215-nm luminescence at room temperature. The intensity of the 215-nm band is uniform over almost all the h-BN surface. In contrast, Cassabois et al. [112] showed that the emission spectrum of h-BN in the deep UV is profoundly structured by phonon-assisted recombination, demonstrating that the bandgap is indirect with a value of 5.955 eV. An exciton binding energy of about 130 meV is derived from the phonon-assisted two-photon absorption spectroscopy. Besides h-BN bulk single crystal, the optical properties of h-BN films prepared by various vapor-phase techniques have also been investigated. For example, a 226-nm CL and an intense exciton-related luminescence at 216–227 nm are observed for the h-BN films grown using the flow rate modulation epitaxy (FME) method and the hBN films synthesized by CVD. As demonstrated by the above results, the direct/indirect nature of the bandgap in h-BN is still controversial. This discrepancy can be understood as follows. First, the LDA-based calculations showed that stacking of BN sheets in the bulk material affects the details of the band structure. Some stacking structures of h-BN are shown to have indirect bandgaps ranging from 2.9 to 4.5 eV while other stacking structures are shown to have direct bandgaps at the K-point. Experimentally, h-BN is likely formed in various stackings, some of which can have direct bandgaps while others are indirect, possibly in the same sample. For example, AB stacking, where every atom in one plane is on top of a different atom in the next plane, widely assumed in theoretical works, is likely an indirect (K–M) bandgap structure. On the other hand, AA stacking, where every atom in one plane is on top of the same atom in the next plane, is predicted by GGA to have a direct bandgap at the K-point. Although AA stacking could be metastable, the cohesive energy difference between the AA and AB structures is only 0.017 eV per atom. Therefore, depending on the specifics of synthesis, either direct or indirect bandgap material could be realized. Moreover, some calculations indicate indirect bandgaps only slightly smaller (50–150 meV) than the direct gaps at the K-point for several polytypes of h-BN. Therefore, h-BN, although technically having an indirect bandgap, can behave as a direct bandgap material in experiments [1].

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4.5.2 Growth techniques for h-BN films 4.5.2.1 Bulk growth of h-BN crystals Balmain first synthesized h-BN in 1842 by reaction of molten boric acid with potassium cyanide. Now, there are mainly two reactions used to produce h-BN on the industrial scale. The first approach to produce h-BN is the reaction of boric acid with ammonia in the presence of carrier substances, for example, CaCO3, Ca3(PO4)2, and BN. The h-BN crystallites synthesized by this method are thin hexagonal platelets with a thickness of about 0.1–0.5 μm and a diameter up to 5 μm. Another way to produce h-BN is the reaction of boric acid or organic nitrogen compounds in nitrogen atmosphere; however, t-BN is usually obtained in this case. Bulk h-BN crystals with larger sizes can be synthesized by the HTHP method by using commercially available h-BN or amorphous BN powders as precursors. Because h-BN is thermodynamically stable at atmospheric pressure, it is also possible to prepare h-BN single crystals at atmospheric pressure and high temperatures. A variety of solvents have been used to produce h-BN crystals, such as Ba3B2N4, Mg3BN3, silicon, sodium, lithium, copper, nickel, NiMo, and Ni-Cr. It is found that alkaline-earth metal solvents with BN (e.g., Ba-BN and MgBN) are useful to synthesize high-purity h-BN crystals under high pressures while they are not applicable for synthesis at atmospheric pressure. On the other hand, transition-metal solvents (Ni, Ni–Mo, Ni–Cr) have been found to be efficient for the recrystallization of h-BN both at high and atmospheric pressures. For example, the synthesis of high-quality h-BN crystals under atmospheric pressure using Ni-Mo as solvent and deoxidized h-BN powder as a starting material was reported by Kubota et al. [113], in which the h-BN crystals with dimensions of several hundred micrometers were obtained (Fig. 4.19B). The h-BN crystals were grown under HTHP via a temperature-gradient method by applying the Ba–B–N solvent system and source of deoxidized h-BN, and the average size of h-BN single crystals is 1 mm2. Recently, transparent and colorless h-BN single crystals were grown from the Mg-B-N system under HTHP [114]. The largest plate-like-shaped h-BN crystals with a size up to 2.5 mm in length and up to 10 μm in thickness were obtained at 30 kbar and 1900–2100°C (Fig. 4.19A). In addition, bulk h-BN crystals 5 mm in size and 40 μm in thickness (Fig. 4.19B) were also synthesized at atmospheric pressure at 1700°C using an Ni-Cr flux [115].

FIG. 4.19 (A) Transparent and colorless h-BN crystals 2.5 mm in size grown from the Mg-B-N system under HTHP. (B) Bulk h-BN crystals 5 mm in size synthesized at atmospheric pressure using an Ni-Cr flux. (A) Reproduced with permission from N.D. Zhigadlo, Crystal growth of hexagonal boron nitride (hBN) from Mg-B-N solvent system under high pressure, J. Cryst. Growth 402 (2014) 308. Copyright 2014, Elsevier.

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4.5.2.2 Growth of h-BN thin film CVD was one of the first techniques used to deposit h-BN thin films. Most of the reports on h-BN deposition during the last years are in combination with the deposition of low-pressure c-BN film, and the crystal structure of the BN is a-BN or t-BN because a t-BN interlayer is usually formed prior to the nucleation of c-BN. An ideal precursor for the CVD growth of h-BN would seem to be borazine (B3N3H6), having the desired 1:1 B/N stoichiometry. However, it has been reported that borazine usually yields t-BN films, which must be annealed to obtain polycrystalline h-BN films with a high amount of t-BN and a-BN. In many CVD processes, a combination of separate precursors is used, and then the ratio of the precursors will be an important process parameter for the CVD growth of h-BN. Ammonia (NH3) is almost always used as a nitrogen precursor, whereas there are several choices for a boron precursor, for example, BF3, BCl3, and BBr3. For CVD processes with halide, it is necessary to remove the halogen atoms from the growing surface, and thus, a high surplus of hydrogen is often needed. To deposit crystalline h-BN by CVD, temperatures above 1100oC and a high N/B ratio are usually necessary. In 2008, thick h-BN epitaxial films were successfully grown on sapphire (0001) by metalorganic vapor phase epitaxy (MOCVD) [116]. They found that the structure of the BN film strongly depends on the V/III ratio and that the h-BN growth under a high V/III ratio can lead to the growth of (0001) h-BN epitaxial films on sapphire substrates. XRD and reflection high-energy electron diffraction (RHEED) reveal that the h-BN epitaxial film has a smooth surface, and the epitaxial relationship is found to be (0001)h-BN jj(0001)sapphire with an in-plane orientation relationship of [11–20]h-BN jj[1100]sapphire [117]. Recently, the epitaxial growth of wafer-scale and wrinkle-free h-BN films on a 2-in. sapphire substrate has been realized by both low-pressure CVD and MOCVD methods [118, 119]. The grown epitaxial layers demonstrate a large absorption coefficient ( 106 cm1) above the bandgap energy of 5.87 eV with direct band transition behavior. Near-band-edge luminescence at 216.5 nm (5.73 eV) and characteristic defect band recombination at longer wavelengths are observed by CL at 77 K. In some cases, a 20-nm BN or AlN buffer layer is first deposited on the sapphire substrate at 800°C, followed by the h-BN growth at high temperatures (1300oC) [120]. Besides sapphire (0001), Ni (111) and 6H-SiC have also been used as a substrate for the epitaxial growth of h-BN. For instance, a dominant band edge emission at around 5.5 eV is observed for h-BN on 6H-SiC substrates, whereas the epitaxial h-BN on Ni (111) exhibits near-band-edge UV luminescence at 227 nm at room temperature [121]. More recently, by using Ni (111) single-crystal as the substrate, centimeter-sized epitaxial h-BN films were obtained by Oh et al. [122] using the atmospheric pressure CVD method with a single ammonia-borane precursor. Only a few reports are available for the growth of h-BN thin films by other methods, such as MBE and PVD. Polycrystalline h-BN films were deposited on sapphire substrates at 850oC by evaporation of elemental boron with NH3 as the nitrogen source, whereas epitaxial growth of a 100-nm-thick h-BN layer was achieved on Ni (111) substrate using an RF nitrogen plasma source in an MBE system. The h-BN layers with superior optical properties in the deep UV were also grown by high-temperature MBE on sapphire (0001) and highly oriented pyrolytic graphite (HOPG) substrates at temperatures between 1200–1700oC [123]. Furthermore, Sutter et al. [124] demonstrated the growth of h-BN films with controlled thickness by magnetron sputtering of a boron target in N2/Ar.

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4.5.2.3 Properties and applications of h-BN film Due to its high chemical and thermal stability and excellent thermal conductivity, h-BN is widely utilized in industry as a solid lubricant, protective coating, insulator, thermal conductor, material for crucibles and linings, microwave-transparent shields, and high temperature bearings. This section focuses on the applications of h-BN as a highly attractive wide bandgap semiconductor material for electronics and optoelectronic devices. So far, fabrication of h-BN-based electronic devices is still in its infant stage of development. The potential of h-BN for deep UV LEDs was recognized more than a decade ago, when Watanabe et al. [111] demonstrated room-temperature UV lasing at 215 nm by accelerated electron excitation for h-BN single crystals obtained by the HTHP method. Later, the obtained h-BN crystals synthesized at atmospheric pressure using an Ni base solvent also emitted intense 215-nm luminescence at room temperature. In 2009, a far-UV plane-emission compact device equipped with a Spindt-type field-emission array as an excitation source was fabricated, in which pure h-BN powder was used as a luminescent material [125]. The device demonstrated stable operation at 225 nm with an output power of 0.2 mW at an accelerating voltage of 8 kV, and the fluctuation in the output was only a few percent. Bulk crystal growth techniques can only produce a millimeter-sized h-BN sample, limiting the prospects for implementing it as an active device material. Recently, Jiang and coworkers demonstrated that the synthesis of wafer-scale semiconducting h-BN epitaxial layers with high crystallinity is possible with MOCVD, thereby providing an opportunity to explore h-BN as an active material for deep UV optoelectronic device applications [120, 126]. The MSM detector consists of microstrip interdigital fingers of Schottky contact formed by a bilayer of Ni/Au [126]. As shown in Fig. 4.20A, the h-BN MSM detectors exhibit a peak E (eV) 3

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(A) The relative spectral response of the h-BN MSM detector measured at 30 V. The inset is a microscope image of the h-BN MSM photodetector with a device size of 1.250.8 mm and 4/4 μm finger width/spacing used for the measurements. (B) Photocurrent decay kinetics of the h-BN MSM detector measured at room temperature under 212-nm irradiation. Reproduced with permission J. Li, S. Majety, R. Dahal, W. P. Zhao, J. Y. Lin, H. X. Jiang, Dielectric strength, optical absorption, and deep ultraviolet detectors of hexagonal boron nitride epilayers, Appl. Phys. Lett. 101 (2012) 171112. Copyright 2012, American Institute of Physics.

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responsivity of 220 nm, a sharp cut-off wavelength around 230 nm, which corresponds well with the band-edge PL emission peak at 5.48 eV (or 227 nm). An outstanding feature observed from h-BN photodetectors is that there are virtually no detectable responses in the long wavelengths measured up to 800 nm. However, because of the poor crystalline quality of h-BN, the observed deep UV to visible rejection ratio in h-BN MSM detectors is still 2–3 orders of magnitude lower than that of AlN-based detectors. In addition, the photocurrent kinetics of the h-BN MSM detector measured at room temperature shown in Fig. 4.20B exhibit no persistent photoconductivity (PPC) effects. Planar MSM devices were also fabricated using h-BN films grown on Si (111) substrates, demonstrating a strong deep UV response [127]. Recently, Li et al. [128] reported the wafer-scale growth of h-BN layers with thicknesses above 1 μm on 2-inch sapphire substrates by MOCVD. MSM device prototypes were fabricated on the h-BN membrane delaminated and separated from the substrate. The prototypes exhibited a low dark current level of around 2 nA under 100 V and a photoconductivity yield of around 100  20% under deep UV illumination at 205 nm. Notably, h-BN-based detectors reported in the literature show very different performances and cannot be compared directly with each other due to the effect of material quality, contact configurations, applied voltages, UV light power densities, and excitation wavelengths. As in the case of c-BN, both p- and n-type conductivities in h-BN have been achieved by in situ and ex situ doping methods. Hall effect measurements reveal that Si-doped h-BN epilayers grown by the MOCVD technique exhibit n-type conduction at 850 K with an inplane resistivity of  12 Ω cm, an electron mobility of  48 cm2/Vs, and a concentration of 1  1016 cm3 [129]. The Mg-doped h-BN epilayers prepared by the same technique exhibit a p-type resistivity around 12 Ω cm at 300 K, which is a reduction of 5–6 orders of magnitude compared to those of Mg-doped AlN [120]. The estimated EA value in h-BN:Mg is around 31 meV based on the temperature-dependent resistivity measurement, and the Hall effect measurements reveal a free hole concentration of 1.1  1018 cm3 and a mobility of 0.5 cm2/Vs. However, the realization of deep UV LEDs with p-n junctions is hampered by difficulties in obtaining successful n-type doping of h-BN. Accordingly, Jiang et al. [130] have proposed the use of highly conductive p-type h-BN as an electron blocking layer (EBL) and a p-type contact layer in AlGaN-based deep UV LEDs, as shown in Fig. 4.21. Calculations of the band edge alignment between h-BN and wurtzite AlN indicate that the h-BN layer should more efficiently block electrons and inject holes into the AlGaN quantum well active region. Simultaneously, as a p-contact layer, h-BN provides reduced contact resistance and increased UV transparency. Nevertheless, despite these potential advantages, the deep UV light emission under current injection has not yet been demonstrated. Due to a large thermal neutron capture cross-section (3840 barns) of the 10B isotope, h-BN also shows potential for detecting fissile material for solid-state thermal neutron detectors. The h-BN growth on the vertical (111) surfaces of parallel trenches fabricated in Si (110) was explored to achieve a thermal neutron detector [127]. The detection efficiency is measured to be 1.1  0.2% for the 0.36 cm2 area device. The detection efficiency for these devices can be enhanced further by improving the crystalline quality of h-BN and by the optimization of the device structure and fabrication processes. Recently, Majety et al. [131] fabricated a

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FIG. 4.21 (A) Schematic diagram of the AlGaN-based deep UV LEDs with a p-type h-BN as an electron blocking and contact layer. (B) The corresponding energy band diagram of the device layer structure.

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B-enriched h-BN thermal neutron detector from a freestanding 43-μm thick h-BN layer prepared by MOCVD growth and subsequent mechanical transfer from a sapphire substrate. The device exhibited a detection efficiency of 51.4%, which is the highest value reported for semiconductor-based neutron detectors to date.

4.6 Synthesis, properties, and applications of 2D h-BN Recently, two-dimensional (2D) materials have received a great deal of attention due to their unique structure, many fascinating properties, and a wide range of technological applications [132–135]. One of the 2D materials is h-BN, which consists of boron and nitrogen atoms arranged in a hexagonal structure such that each boron atom is bonded to three neighboring nitrogen atoms and vice versa. The resulting structure resembles graphene, with a sim˚ . Because of its analogous crystal structure and small lattice ilar lattice constant of 2.52 A mismatch (1.7%) with graphene, atomically thin h-BN has many properties similar to those of graphene. For example, the calculated Young’s modulus of monolayer h-BN is 0.71– 0.97 TPa and the breaking strength is 120–165 GPa, whereas the thermal conductivity of few-layer h-BN is measured to be 100–270 Wm1 K1. On the other hand, unlike semimetal graphene, h-BN is an insulator with a wide bandgap (5.97 eV). Due to the atomically smooth surface that is free of dangling bonds and with negligible defects as well as only a low dielectric constant and high breakdown field, h-BN layers can be used as dielectric substrates for graphene- and transition metal dichalcogenide (TMD)-based heterostructures for the next generation of electronic and optical devices. Indeed, the mobility for graphene devices on h-BN was observed to be an order of magnitude higher than that of the devices fabricated on the commonly used SiO2/Si substrate, and micrometer-scale ballistic transport and a

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quantum Hall effect have been observed in the h-BN-based 2D heterostructures. Furthermore, owning to its wide bandgap, h-BN can also be applied as a barrier layer to fabricate the tunneling devices and tune the carrier dynamics and optical properties. Additionally, the characteristic resistance to oxidation and corrosion makes h-BN a suitable candidate as a gate dielectric and capping layer to protect the active material element and device from oxidation and chemical degradation. For electronic applications, high-quality and large-area h-BN layers with few defects are strongly desirable. During the past decade, much effort was directed toward producing mono- and few-layer h-BN sheets. Similar to graphene, h-BN flakes can be produced by mechanical or liquid-phase exfoliation; nevertheless, the limited flake size hinders their application in large-area devices. Monolayer and few-layer h-BN have been grown on single-crystal transition metals such as Ru (001), Rh (111), and Pt (111) with expensive ultrahigh vacuum (UHV) systems by CVD from borazine or magnetron sputtering of the boron target in Ar/N2, respectively. Despite the great progress achieved so far, the current research into 2D h-BN still faces three major challenges: (1) growth of large-scale singlecrystal h-BN layers with controlled layer numbers; (2) modification of h-BN via doping and functionalization, or production of hybridized nanosheets containing h-BN phases; and (3) integration of h-BN into other nanodevices. In this part, we will first describe the recent progress in the synthesis of 2D h-BN, including exfoliation, CVD, ion beam sputtering deposition (IBSD), cosegregation, and several other new fabrication techniques, together with their growth features. Then, we will focus on the intriguing properties as well as the potential applications of the h-BN layer and related heterostructures.

4.6.1 Synthesis techniques of 2D h-BN The typical synthesis strategies of 2D h-BN are divided into two major routes: top down and bottom up. The top-down approach for the synthesis of 2D h-BN mainly includes mechanical exfoliation and chemical exfoliation. The difference between the two methods lies in the way in which the van der Waals forces between the atomic layers are broken. Bottom-up routes, in which atoms are assembled in order on specific substrates through physical or chemical methods, offer significant advantages for synthesizing large-area 2D h-BN compared to the top-down method. Among the bottom-up methods, CVD is the most thoroughly explored and promising for the large-scale production of 2D h-BN, which is a key prerequisite for its applications. Besides CVD, other novel methods such as IBSD, sputtering, and cosegregation have also been developed during the past several years for synthesizing 2D h-BN. 4.6.1.1 Mechanical exfoliation Mechanical exfoliation technique by Scotch was initially used to isolate graphene in 2004. Due to the weak van der Waals bonds between adjacent layers, similar to graphene, h-BN nanosheets can be easily produced by mechanical exfoliation [136]. The exfoliation technique yields high-quality monolayers with fewer defects than those produced by chemical methods. Therefore, 2D h-BN layers produced by the exfoliation method are used for exploring their intrinsic properties or fundamental research in electronics and optoelectronics. Nevertheless, the h-BN layers obtained by mechanical exfoliation are often randomly

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distributed with limited flake size and extremely low yield, partially due to the stronger interactions between BN basal planes. The overall sizes of the exfoliated h-BN nanosheets did not usually exceed several micrometers due to the size limitation of pristine h-BN powders. Recently, the few-layer h-BN samples with a size of more than 100 μm were able to be obtained by mechanical exfoliation employing large h-BN single crystals [137]. In addition to direct mechanical exfoliation, Li et al. [138] have demonstrated a low-energy method, that is, a controlled ball milling process, to produce high-quality h-BN in high yield and efficiency with little damage to the layer structure. To improve the production of h-BN nanosheets, a low-power sonication process is combined with the low-energy ball milling approach. The mechanical exfoliation method is the simplest route to obtain h-BN nanosheets with high crystal quality. Due to its simplicity and operability, the mechanical exfoliation process is usually combined with the sequential transfer process or CVD of the other layered materials. Two or more different 2D materials can be assembled together and form vertical heterostructures based on h-BN nanosheets. However, the limited flake size hinders their application in large-area devices.

4.6.1.2 Liquid-phase exfoliation Liquid-phase exfoliation, which was first reported in 2008 by Han et al. [139], creates dispersions of 2D h-BN in various solvents or aqueous surfactant solutions with the assistance of sonication. Sonication of sufficient intensity breaks the van der Waals bonds between the atomic layers of h-BN. This allows the solvent molecules to seep between the boron nitride layers and expand, which results in complete exfoliation of h-BN into mono- or few-layers, forming a dispersion of h-BN nanosheets. Due to the stronger interlayer force of h-BN, the solvents used for the exfoliation of h-BN are usually strong solvents, such as the strong-polar N,N-dimethylformamide (DMF), Lewis bases Octadecylamine (ODA), polyethylene glycol (PEG), methane sulfonic acid, etc. Lin et al. [140] showed that by simply utilizing water as the solvent, few-layered h-BN could be obtained because of the hydrolysis of h-BN under sonication participates. The critical point of liquid-phase exfoliation is to reduce the mixing enthalpy by choosing the correct solvents whose surface tension matches with the surface energy of h-BN. Zhou et al. [141] demonstrated a versatile and scalable mixed solvent strategy for liquid exfoliation of h-BN. Besides, there are several other strategies based on liquid-phase exfoliation that have been explored, such as electric field-assisted liquid exfoliation, microwave irradiation-assisted exfoliation, liquid exfoliation combined with fast quenching, magnetic stirring-assisted ultrasonication, etc. Liquid-phase exfoliation is quick and easy and insensitive to ambient conditions, leading to higher yields. It also exhibits the ability to produce thin films of inorganic layered compounds with applications in batteries, coatings, and lubricants. Moreover, the liquid-phase exfoliation method enables the production of a wide range of hybrids with adjustable thermal and electrical conductivity and attractive thermoelectric properties as well as composites with enhanced mechanical properties. However, this method suffers from surface contamination of the samples and the limited flake size, and the number of layers and their lateral size are difficult to control in contrast with the other techniques.

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4.6.1.3 Chemical vapor deposition CVD has been applied to epitaxially grown h-BN thin films for several decades. The common strategy for the CVD growth of single or a few layers of h-BN is decomposition of the source precursors, such as borazane (BH3NH3) or benzene-like borazine (B3N3H6) and trichloroborazine (B3Cl3H3N3). In the early period, the h-BN monolayer was epitaxially grown on single-crystal transition metals such as Pt (111), Ru (0001), and Ni (111) with UHV systems by CVD through the decomposition of B3N3H6. The transition-metal substrate materials have a catalytic effect and assist in decomposition of the precursors and formation of monolayer h-BN. Once the substrate is covered with h-BN, no further decomposition of the precursor occurs, and the growth becomes self-limited. Thus, it is rather difficult to control the number of layers of h-BN. In addition, the UHV process requires expensive and sophisticated equipment, and these studies were only focused on the electronic and surface properties by various in situ characterization techniques. Recently, both atmospheric-pressure CVD (APCVD) and low-pressure CVD (LPCVD) have been used to grow 2D h-BN on polycrystalline metallic substrates [142–145]. Fig. 4.22A schematically shows the experimental setup for the CVD growth of h-BN with the Cu catalytic substrate [144]. As shown in Fig. 4.22B, the CVD process involves four distinct steps: the adsorption of precursor molecules on the substrate surface, the decomposition of the precursor and formation of mobile surface species, the diffusion of these species, and nucleation and/or incorporation into the growing film. The growth of h-BN depends on various factors such as precursors, substrates, growth temperature and pressure, and hydrogen partial pressure. To improve the quality and functionalities of h-BN films, intensive research efforts have been paid to enlarge the grain size, fabricate the h-BN film of desired layer numbers, and precisely control the grain morphology and orientation. Some important growth parameters and aspects will be described as follows. Various gaseous, liquid, or solid precursors are widely utilized in CVD processes. Gaseous precursors such as BF3/NH3, BCl3/NH3, and B2H6/NH3 have been used to deposit 2D h-BN.

FIG. 4.22 (A) Schematic illustration of the CVD experimental setup with Cu foils in a ceramic crucible. (B) Schematic illustration of the mechanisms in the growth of h-BN on the Cu surface by CVD. Reproduced with permission R.-J. Chang, X. Wang, S. Wang, Y. Sheng, B. Porter, H. Bhaskaran, J. H. Warner, Growth of large single-crystalline monolayer hexagonal boron nitride by oxide-assisted chemical vapor deposition, Chem. Mater. 29 (2017) 6252. Copyright 2017, American Chemical Society.

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For gaseous precursors, the ratio between the boron source and ammonia is the critical factor for preparing stoichiometric h-BN. However, their use has been gradually reduced because of their toxicity. Unlike gaseous precursors, liquid single-compound precursors such as borazine (B3N3H6), trichloroborazine (B3N3H3Cl3), or hexachloroborazine (B3N3Cl6) have an advantage of 1:1 B/N stoichiometry. Especially, the decomposition of borazine does not produce the high toxic byproducts such as BF3 or BCl3. But borazine is sensitive to moisture and easily hydrolyzes to boric acid, ammonia, and hydrogen in water. Borazane (or ammonia borane) is a crystalline solid at room temperature and melts at around 106oC. It not only has 1:1 B/N stoichiometry but is easily accessible and more stable under ambient conditions. The decomposition of borazane produces hydrogen, monomeric aminoborane (BH2NH2), and borazine. In order to use it as a precursor for the CVD growth of h-BN, the ammonia borane needs to be heated to generate borazane vapor and to diffuse onto the substrates. It is found that the ambient pressure has an important effect on the growth of 2D h-BN layers [146]. The lower pressure usually leads to h-BN that is uniform in thickness, highly crystalline, and consists solely of h-BN. Although the larger pressure leads to the increase of the growth rate, the resulting h-BN is more amorphous, disordered, and sp3-bonded. This phenomenon has been attributed to the incomplete thermolysis of the H3N-BH3 precursor from a passivated Cu catalyst under high pressure. Furthermore, hydrogen added to the CVD precursor gases has strong effects in the morphology of the growing h-BN domains. As catalyst-type substrates, polycrystalline Cu and Ni foils are traditionally used for the CVD growth of h-BN due to their low cost and excellent catalytic performance. The morphology and nucleation density are highly dependent on the pretreatment of the substrates, and a short acetic acid wash has been found to be effective to remove the oxidized layer and contamination from the Cu surface. On the polycrystalline Cu substrates, surface irregularities such as grain boundaries and the rough surface of Cu are known to stimulate h-BN nucleation. Electropolishing and annealing at high temperature can minimize the defects on the Cu surface, which is favorable for the growth of h-BN domains with larger lateral sizes. Comparable to solid metal substrates, the surface of liquid metals is more uniform with fewer grain boundaries. The homogenous nucleating and orderly arranged h-BN layer was achieved on the melt Cu substrates by APCVD, and the h-BN layers grown on liquid Cu substrates are self-aligned single-crystal arrays. Besides Cu and Ni, other metallic substrates such as Pt, Co, Cu-Ni, and Ni-Ga alloy have also been utilized for the growth of 2D h-BN. In addition to the commonly used metal substrates, other catalyst-free dielectric substrates such as Ge, sapphire, Si, and SiO2-coated Si were also explored. It was reported that few- to multilayer continuous nanocrystalline h-BN films with no observable pinholes or wrinkles were achieved directly on both amorphous SiO2/Si and quartz substrates without any metal catalyst. For example, Behura et al. [147] prepared h-BN film directly on SiO2 and quartz substrates, where the oxygen on these surfaces binds with boron to produce nucleation sites for oxide-assisted growth of large-area and continuous h-BN films. Recently, Lee et al. [148] devised a promising strategy for the growth of wafer-scale and defect-free h-BN on sapphire substrates, where nanocrystalline graphene acted as a seed layer for the h-BN growth. By using this method, a wafer-scale ultraflat h-BN thin film was successfully synthesized on a bare sapphire substrate. CVD is probably the most likely method to develop into a technique used for the mass production of high-quality 2D h-BN. A practical problem during the CVD growth of h-BN is the use of some unconventional precursors, as mentioned above. These compounds are rare and

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highly toxic, unstable, or pyrophoric. In addition, the CVD growth relies on the decomposition of precursors, the rate of which may change dramatically with the exposed surface area of the catalytic substrate. 4.6.1.4 Physical vapor deposition (PVD) PVD is a common process for film growth that can avoid the complex interrelation in growth parameters involved in the CVD process. In view of this point, Sutter et al. [124] deposited mono- and few-layer h-BN films with controlled thickness on an Ru (0001) surface by RF magnetron sputtering of a boron target under N2/Ar atmosphere (Fig. 4.23A). The h-BN films up to two atomic layers can be synthesized by reactive deposition at high substrate temperatures. Whereas direct growth onto clean Ru (0001) produces h-BN (A  A0 A stacking, top image in Fig. 4.23B), the termination of the Ru surface with a monolayer graphene (MLG) prior to BN growth induces a different stacking order (A  B  C, bottom image in Fig. 4.23B), that is, produces r-BN.

FIG. 4.23 (A) Schematic diagram of RF magnetron sputtering for few-layer h-BN deposition from a solid boron target in Ar/N2 gas. (B) HRTEM image of a three-layer h-BN film on Ru (0001) (top image) and a 10-layer A BC stacked r-BN on graphene-covered Ru (0001) (bottom image). (C) Schematic diagram of the IBSD system in which boron and nitrogen species were sputtered from an h-BN target by Ar ions with a Kaufmann ion source. SEM images of h-BN domains deposited on the Cu foils at 1050°C for 4 min under (D) pure Ar and (E) Ar/H2 mixed atmosphere. (A) and (B) Reproduced with permission from P. Sutter, J. Lahiri, P. Zahl, B. Wang, E. Sutter, Scalable synthesis of uniform fewlayer hexagonal boron nitride dielectric films, Nano Lett. 13 (2013) 276. Copyright 2013, American Chemical Society. (C)–(E) Reproduced with permission from H.L. Wang, X.W. Zhang, J.H. Meng, Z.G. Yin, X. Liu, Y.J. Zhao, L.Q. Zhang, Controlled growth of few-layer hexagonal boron nitride on copper foils using ion beam sputtering deposition, Small 11 (2015) 1542. Copyright 2015, John Wiley and Sons.

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Wang et al. [149, 150] reported a facile and innovative method for growing high-quality few-layer h-BN by IBSD. Compared to magnetron sputtering, one benefit of IBSD is that the Ar plasma is restricted in the ion source, eliminating the side effect of plasma in the sputtering process. Prior to the IBSD growth, the Cu or Ni foils were annealed at 1000oC for 20 min under a hydrogen atmosphere, and then boron and nitrogen species were sputtered from an h-BN target by 1.0 keV Ar ions with a Kaufmann ion source, resulting in the growth of 2D h-BN, as shown in Fig. 4.23C. The domain density of h-BN can be reduced by introducing hydrogen into the deposition chamber, resulting in an increased size of h-BN domains (Fig. 4.23D and E). Moreover, they also reported the deposition of hBN on graphene by IBSD [151]. The bombardment of energetic particles during the deposition of h-BN leads to the mixing of boron/nitrogen atoms and carbon atoms, and as a result, the atomic layer of hybridized h-BNC was synthesized. Other PVD-based strategies such as MBE and PLD have also been exploited to prepare h-BN with controlled numbers of layers. Compared to the conventional CVD method, the PVD technique avoids the use of unconventional precursors and is much easier to control, which should be very useful for the largescale production of 2D h-BN layers. In addition, it is possible to exploit this approach to directly grow h-BN on other substrates such as SiO2/Si and sapphire. There is no doubt that these PVD techniques provide a different approach toward the high-quality and wafer-scale fabrication of 2D h-BN, although the crystallinity of the h-BN films needs to be optimized while complex fabrication systems are also required. 4.6.1.5 Cosegregation method Surface segregation appears to be another feasible method for the large-scale synthesis of 2D layered materials. It was first used by Xu et al. [152] to synthesize 2D h-BN. They reported the formation of monolayer and few-layer h-BN by segregation of boron and nitrogen atoms from a bulk Fe–Cr–Ni alloy doped with boron and nitrogen. Suzuki et al. [153] have also grown fewlayer h-BN thin films by annealing the Co(Ni)/amorphous h-BN/SiO2 structure, where h-BN was embedded between the polycrystalline metal and the SiO2 substrate, as shown in Fig. 4.24A. The h-BN film thus formed on the topmost surface via diffusion of boron and

FIG. 4.24 Schematic of the cosegregation method. (A) Schematic of the growth method for a few-layer h-BN by annealing the Co(Ni)/amorphous h-BN/SiO2 structure h-BN film in vacuum. (B) Schematic of h-BN synthesis by the vacuum annealing of sandwiched substrates Fe/(B,N)/Ni.

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nitrogen atoms through the metal films. To better control the segregation process, Zhang et al. [154] designed a sandwiched substrate with a solid-state (B,N)-layer between an Fe top layer and an Ni bottom layer, as illustrated in Fig. 4.24B. This method follows an underneath-growth model, and patterned h-BN thin films are achieved by prepatterning the solid (B,N)-source. Moreover, the h-BN thin films were also prepared by utilizing diffusion and surface segregation of boron and nitrogen in Ni and Co thin films on SiO2/Si substrates after exposure to diborane and ammonia precursors at high temperature [155]. It was found that the cooling rates have no effect on the segregation kinetics in thin films due to the high boron and nitrogen diffusivities. The resulting h-BN films have a near theoretical optical bandgap and excellent electrical breakdown strength. The segregation-precipitation process provides a unique approach for the deposition of h-BN thin films on arbitrary substrates and can be extended to wafer scale. Other processes such as h-BN nanotube unwrapping and the substitution reaction of graphene produce 2D h-BN that are not high purity or large in size. Although several synthesis routes have been identified to yield mono- and few-layer h-BN, the production of large and single-crystalline h-BN domains with controlled thicknesses for practical use is still a challenge. Variable thicknesses and impurities are the main drawbacks of the wet chemical routes. As far as h-BN as a substrate for 2D nanoscale electronics is concerned, among the various synthesis routes, the CVD method has been deemed to be an effective route to produce large h-BN domains for practical applications. The major challenges for CVD growth are how to minimize the grain boundaries and how to control the formation of mono- or fewlayer h-BN over a large area.

4.6.2 Growth features of 2D h-BN CVD-grown h-BN single-crystal domains on metal substrates usually adopt a triangle shape instead of the hexagonal shape observed for graphene. Theoretical calculations also suggest that the most favorable configuration of the triangular h-BN crystal is with an N-terminated edge in a zigzag direction under a hydrogen atmosphere [156]. Besides h-BN triangles, other shapes such as hexagons, trapezoids, asymmetric diamonds, and irregular circles are also observed. Stehle et al. [157] found that h-BN crystals change their shape from triangular to truncated triangular and further to hexagonal, depending on the Cu substrate distance from the precursor (Fig. 4.25). Such evolution in shape is believed to be associated with increasing the boron to nitrogen ratio along the CVD tube and the chemical potential of feeding species. Using a wedge-shaped Cu pocket with two opened ends, Yin et al. [158] demonstrated that h-BN triangles grown on the inside of the pocket gradually evolve to truncated triangles and then to hexagons with decreasing the gap distance between the top and bottom Cu foil. This kind of substrate position effect is also observed on a resolidified Cu surface. Several theoretical works have been conducted to explain the morphology of h-BN domains. Liu et al. [156] found that depending on the chemical potential of constituent elements, h-BN shapes as equilateral triangles or hexagons can be formed. Zhang and coworkers [159] revealed an atomistic growth mechanism of h-BN based on nonequilibrium dynamics near the growing edge. For Cu foils at N-rich condition, the h-BN islands are N-terminated triangles; under B-rich condition, the BN shapes can evolve into truncated triangles or hexagons with additional B-terminated edges.

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FIG. 4.25 APCVD experimental setup for h-BN growth. Copper substrates were placed at the positions A, B, and C with an 8-in. distance between them. The corresponding SEM images of the h-BN domains grown at 1065°C at the positions A, B, and C using argon as a buffer gas. Sketch of the resulting h-BN crystal shapes and corresponding termination-nitrogen (blue) and boron (red). Reproduced with permission from Y. Stehle, H.M. Meyer, R.R. Unocic, M. Kidder, G. Polizos, P.G. Datskos, R. Jackson, S.N. Smirnov, I. V. Vlassiouk, Synthesis of hexagonal boron nitride monolayer: control of nucleation and crystal morphology, Chem. Mater. 27 (2015) 8041 Copyright 2015, American Chemical Society.

The area covered by the h-BN monolayer and the crystal size is a very important consideration for the practical application of 2D h-BN films. Generally, large-size domains lead to a low density of domain boundary defects, resulting in a higher quality of film while the limited domain size and unavoidable domain boundaries in h-BN thin films always significantly impair the performance of electronic devices. In earlier reports, the average sizes of CVD-grown h-BN domains were limited to several micrometers [132–135]. Thus far, numerous efforts have been devoted to improve the domain size of h-BN. It is known that the rough surface or presence of grain boundaries is likely to act as a nucleation seed, leading to an enhanced rate of nucleation. Thus, to obtain single-crystal h-BN domains that are as large as possible, a common strategy is reducing the nucleation density during the growth process. Pretreatments of the metal substrates, including electrochemical polishing and annealing for hours at high temperature, are imperative procedures to decrease the surface

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roughness and increase the domain sizes. Tay et al. [160] demonstrated the growth of hexagon-shaped h-BN domains with an area of 35 μm2 on highly smoothed electropolished Cu substrates by APCVD (Fig. 4.26A). The nucleation density of h-BN grains can also be reduced by increasing the thermal annealing duration of the Cu foil and the size of the produced triangle-shaped h-BN domains can be increased to 20 μm [161]. On the other hand, reducing the precursor feeding rate proves to be the other route to increase the domain size of h-BN grown on Cu substrates. Song et al. [162] used a folded Cu enclosure approach to suppress the precursor feeding rate and improve the surface smoothness, and the resulting h-BN domain size is up to 72 μm in edge length (Fig. 4.26B). It also was reported that the density of the nucleation seeds can be decreased with increasing the hydrogen gas flow, leading to the formation of a 100-μm single-crystal h-BN domain (Fig. 4.26C) [163]. Due to the high chemical affinity of N-containing intermediate species to the Cu surface, the nucleation density of h-BN on Cu foils is usually higher, resulting in the limited-size h-BN domains. Thus, a special design of the substrate is required for further enlarging the h-BN domain size. By using the Ni foil substrates combined with the in situ pretreatment to eliminate surface irregularities and remove impurities, Wang et al. [150] have grown large single-crystal h-BN domains with a lateral size up to 110 μm on Ni foils using IBSD (Fig. 4.26D). Similarly, by choosing the Cu–Ni alloy as the substrate, Lu et al. [164] developed single-crystal monolayer h-BN domains with the length of the side up to 130 μm. The introduction of Ni could effectively reduce the nucleation density to 60 per mm2, thus giving rise to large sizes of h-BN domains (Fig. 4.26E). Likewise, an active Si-doped Fe substrate could be applied to effectively tune the nucleation density of h-BN via controlling the amount of Si diffusion into the Fe thin film. As a result, Caneva et al. [165] synthesized the hBN domains with lateral dimensions of 300 μm and continuous h-BN monolayer films with large domain sizes (25 μm) (Fig. 4.26F). To further enhance the size of the h-BN domain, a combined strategy including a Cu enclosure made from a thick Cu foil (127 μm), highpressure preannealing, and an inserted zigzag bent copper foil has been adopted in an LPCVD, which has led to successfully obtaining single-crystal monolayer h-BN domains with a lateral size up to  300 μm (Fig. 4.26G) [166]. Very recently, Meng et al. [167] demonstrated the growth of millimeter-sized single-crystal h-BN domains on epitaxial Ni (111)/ sapphire substrates by the IBSD technique. They found that the Ni (111) film with a smooth surface and less grain boundary is a critical factor for growing large-size h-BN domains. Under the optimized growth conditions, single-crystal h-BN domains up to 0.6 mm in edge length can be obtained (Fig. 4.26H). As discussed above, a number of intensive efforts have been made to grow an h-BN monolayer with a large domain size via CVD. Nevertheless, further enlarging the grain size requires a long growth time and thus becomes more difficult. The second approach for reducing or even eliminating domain boundaries is growing aligned h-BN domains and stitching them together to form a uniform h-BN layer. It is expected that the unidirectionally aligned h-BN nuclei can grow and then coalesce into a single crystal without grain boundaries, similar to graphene. Actually, it has been reported that aligned monolayer h-BN domains are epitaxially grown on single-crystal substrates such as Ru (0001), Ir (111), and Ni (111) in the UHV system [168, 169]. However, these studies are only focused on the electronic and surface properties by various in situ characterization techniques, and noble metal substrates are not preferred for large-scale application.

FIG. 4.26 Large-sized h-BN domains. (A) Hexagonal-shape h-BN domains with an area of up to 35 μm2 obtained on highly electropolished Cu substrates. (B) A large-domain h-BN triangle showing an edge length of 72 μm achieved by reducing the precursor feeding rate. (C) Formation of a 100-μm single crystal h-BN domain by increasing the hydrogen gas flow. (D) Single-crystal h-BN domains with a lateral size up to 110 μm obtained by the in situ pretreatment of Ni substrates. (E) Single-crystal h-BN domains with the length of side up to 130 μm achieved by choosing the Cu–Ni alloy as the substrate. (F) The h-BN domains with lateral dimensions of 300 μm obtained by controlling the amount of Si diffusion into the Fe substrate. (G) Single-crystal h-BN domains with a lateral size up to  300 μm obtained by a combined strategy, including a Cu enclosure made from a thick Cu foil (127 μm), high-pressure preannealing, and an inserted zigzag bent copper foil. (H) Single-crystal h-BN domains up to 0.6 mm in edge length synthesized on epitaxial Ni (111)/sapphire substrates by IBSD. (A) Reproduced with permission from R.Y. Tay, M.H. Griep, G. Mallick, S.H. Tsang, R.S. Singh, T. Tumlin, E.H.T. Teo, S.P. Karna, Growth of large single-crystalline two-dimensional boron nitride hexagons on electropolished copper, Nano Lett. 14 (2014) 839. Copyright 2014, American Chemical Society. (B) Reproduced with permission from X.J. Song, J.F. Gao, Y.F. € Nie, T. Gao, J.Y. Sun, D.L. Ma, Q.C. Li, Y.B. Chen, C.H. Jin, A. Bachmatiuk, M. H. Rummeli, F. Ding, Y.F. Zhang, Z.F. Liu, Chemical vapor deposition growth of large-scale hexagonal boron nitride with controllable orientation, Nano Res. 8 (2015) 3164. Copyright 2015, Springer. (C) Reproduced with permission from Q. Wu, J.-H. Park, S. Park, S.J. Jung, H. Suh, N. Park, W. Wongwiriyapan, S. Lee, Y.H. Lee, Y.J. Song, Single crystalline film of hexagonal boron nitride atomic monolayer by controlling nucleation seeds and domains, Sci. Rep. 5 (2015) 16159. Copyright 2015, Nature Publishing Group. (D) Reproduced with permission from H.L. Wang, X.W. Zhang, H. Liu, Z.G. Yin, J.H. Meng, J. Xia, X.-M. Meng, J.L. Wu, J.B. You, Synthesis of largesized single-crystal hexagonal boron nitride domains on nickel foils by ion beam sputtering deposition, Adv. Mater. 27 (2015) 8109. Copyright 2015, John Wiley and Sons. (E) Reproduced with permission from G. Lu, T. Wu, Q. Yuan, H. Wang, H. Wang, F. Ding, X. Xie, M. Jiang, Synthesis of large single-crystal hexagonal boron nitride grains on Cu-Ni alloy, Nat. Commun. 6 (2015) 6160. Copyright 2015, Nature Publishing Group. (F) Reproduced with permission from S. Caneva, R.S. Weatherup, B.C. Bayer, B. Brennan, S.J. Spencer, K. Mingard, A. Cabrero-Vilatela, C. Baehtz, A. J. Pollard, S. Hofmann, Nucleation control for large, single crystalline domains of monolayer hexagonal boron nitride via Si-doped Fe catalysts, Nano Lett. 15 (2015) 1867. Copyright 2015, American Chemical Society. (G) Reproduced with permission from Y. Ji, B. Calderon, Y. Han, P. Cueva, N.R. Jungwirth, H.A. Alsalman, J. Hwang, G.D. Fuchs, D.A. Muller, M.G. Spencer, Chemical vapor deposition growth of large single-crystal mono-, bi-, tri-layer hexagonal boron nitride and their interlayer stacking, ACS Nano 11 (2017) 12057. Copyright 2017, American Chemical Society. (H) Reproduced with permission from J. H. Meng, X.W. Zhang, Y. Wang, Z.G. Yin, H. Liu, J. Xia, H.L. Wang, J.B. You, P. Jin, D.G. Wang, X.-M. Meng, Aligned growth of millimeter-size hexagonal boron nitride singlecrystal domains on epitaxial nickel thin film, Small 13 (2017) 1604179. Copyright 2017, John Wiley and Sons.

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Many efforts have been devoted to growing aligned h-BN domains on various substrates. The Cu (111) facet possesses the same hexagonal symmetry as h-BN with a mismatch of about 1.8%, which is one of the most promising substrates for growing aligned h-BN domains and even singlecrystal h-BN layers. As in the case of nonpolar-polar epitaxy, two opposite lattice orientations of hBN triangles are favored on the Cu (111) facet. In contrast, for h-BN triangles grown on Cu (102) and (103) facets, only one unique orientation is observed [170]. Yin et al. [171] reported that highly aligned h-BN domains with two and four primary orientations can be epitaxially grown on semiconducting Ge substrates of (110) and (100) facets, respectively. Similarly, the aligned growth of single-crystal h-BN domains with antiparallel orientation was realized on the Ni(111)/sapphire substrates by the IBSD technique [167]. Suppressing the oppositely orientated h-BN domains during the growth is a great challenge due to the bipolar structure. It has been proven that h-BN domains with the same orientation can merge seamlessly without structural domain boundaries, making it promising for the fabrication of wafer-scale single-crystal h-BN films.

4.6.3 Properties and applications of 2D h-BN As an inheritor of bulk h-BN, 2D h-BN also exhibits a unique combination of advantageous properties, including a wide optical bandgap, excellent electrical insulation, low dielectric constant, high thermal conductivity, high elastic modulus, excellent inertness, and low friction coefficient. These unique properties make 2D h-BN a promising candidate for a rich variety of applications. The minimal lattice mismatch (1.7%) with graphene and almost atomically flat surface that can be free of dangling bonds or surface charge traps make h-BN a very attractive choice as dielectric layers and insulating substrates for graphene-based devices. Generally, the carrier mobility of graphene on SiO2 is limited by scattering from charged surface states or impurities, substrate surface roughness, and surface optical phonons. It has been experimentally shown that the use of h-BN as support and dielectric layers in graphene FET instead of a conventional SiO2 layer leads to a significant improvement of device characteristics [172–174]. In graphene on h-BN, high mobilities of 125,000 cm2 V1 s1 at room temperature and 275,000 cm2 V1 s1 at 4.2 K, which are close to that of suspended graphene, have been achieved [175]. Wang et al. [176] observed even higher roomtemperature mobility up to 140,000 cm2 V1 s1 in an h-BN/graphene/h-BN device fabricated by the van der Waals assembly technique with edge contacts, virtually reaching the theoretical limit. Fig. 4.27 compares the mobilities of exfoliated monolayer graphene supported on several typical substrates at low temperature (<20 K), highlighting the outstanding performance of h-BN in this aspect [135]. In addition to graphene, other 2D materials such as black phosphorus and TMDs supported on or encapsulated between h-BN sheets also show intrinsic outstanding performances. For example, such improvements were observed in WSe2- and MoS2-based transistors encapsulated by h-BN nanosheets, where a record Hall mobility (>600 cm2 V1 s1) and low subthreshold swing (80 mV/ dec) were achieved, respectively [177, 178]. Due to the weak van der Waals interactions between the atomic layers, h-BN not only acts as a buffer layer for a nitride device but also can form a release layer that enables the mechanical transfer of GaN-based device structures onto foreign substrates in a highly scalable way [179].

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FIG. 4.27 Carrier mobility of exfoliated monolayer graphene supported on several typical substrates at low temperatures (1.7–20 K).

The stacking of various 2D materials in the vertical and lateral directions has been studied recently. With appropriate design, vertical heterostructures combining diverse 2D materials may exhibit surpassing or complementary electronic characteristics that provide further possibilities for electronic device architectures and largely expand the potential applications of 2D materials. The methods employed for forming vertical heterostructures based on hBN can be generally categorized into two groups: (a) transfer of one type of exfoliated or pregrown 2D nanosheet onto another type [174, 177], and (b) direct growth of one type of 2D nanosheets on another type [180–182]. However, the mechanical transfer process for the fabrication of heterostructures is too complicated, inefficient, and nonscalable for real applications. Therefore, a major challenge in future research is to grow heterostructures in large area by CVD without any transfer. From a scientific and technical point of view, direct synthesis of multilayer 2D vertical heterostructures with controlled thickness would be one of the ultimate goals. The absence of an energy gap limits the on/off ratio in graphene transistors, rendering them inadequate for implementation in integrated circuits. To overcome this limitation, Britnell et al. [183] reported a field effect tunneling transistor (FETT) based on vertical graphene/h-BN/graphene heterostructures, in which the atomically thin h-BN layer acts as a barrier layer between two graphene layers. By applying a gate voltage between the Si substrate and the bottom graphene layer, the Fermi level in the graphene layers could be tuned, resulting in the tenability of the tunneling current. An on/off ratio of 50 is obtained at room temperature, which exceeds those demonstrated for planar graphene FETs at room temperature by a factor of 10. Moreover, with bilayer graphene on the h-BN substrate, a very high current-gain cut-off frequency of ft ¼ 33 GHz was demonstrated, which is two times larger than those achieved on SiO2 substrates (18 GHz) and Al2O3 gate dielectrics [174]. A few-layer h-BN introduced between the Si and graphene has been shown to improve the performance of graphene-on-silicon Schottky junction solar cells [184]. In this case, hBN acts as an effective electron-blocking/hole-transporting layer due to its unique properties and appropriate band alignment with Si, and thus the interface recombination is suppressed and the open circuit voltage is remarkably increased. On the other hand,

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the use of the h-BN layer improves the conductivity of graphene and thus the short circuit current through reduced series resistance. Moreover, the interface defects and contamination arising from the layer-by-layer transfer process can be eliminated by using a directly grown graphene/h-BN heterostructure. As a result, a power conversion efficiency of 10.93% has been achieved, which corresponds to an improvement of 15% compared to a reference device based on direct application of graphene contact to Si. Apart from the vertical stacking, in-plane (lateral) graphene/h-BN heterostructures have also attracted much attention in recent years because of the interesting functions when two domains are seamlessly stitched together. A graphene/h-BN lateral structure with atomic lattice coherence can be formed by either the direct growth of h-BN at the edge of graphene, or by chemically converting graphene to h-BN [185–188]. Electronic devices based on graphene/h-BN lateral structures have also been investigated such as FETs, resonators, etc. By changing the concentration of carbon in the channel, the mobility and the on/off ratio of the device can be tuned. Because of the insulating properties of h-BN, electrically isolated graphene devices can be fabricated in a single atomically flat sheet based on the graphene/h-BN lateral heterostructure, as shown in Fig. 4.28A and B [185]. Such hybrid structures are particularly useful for ultraflat three-dimensional electronics. Owing to the short history of lateral heterostructures, many properties of these interfaces are still unknown, but there has been a meaningful advance in the measurement of such properties in recent years. As a wide bandgap semiconductor with UV luminescence, h-BN could be a promising candidate for application in photon emission, deep UV lasing, and detectors. Actually, deep UV photodetectors based on the transferred h-BN layers have been demonstrated (Fig. 4.28C–E) [150, 189]. The h-BN photodetectors exhibited high performance with an on/off ratio of >103 under deep UV light illumination at 212 nm and a cut-off wavelength at around 225 nm with virtually no responses for below bandgap excitation [189]. This work demonstrates that twodimensional h-BN layers are promising for the construction of high-performance solar-blind photodetectors. Moreover, ultrabright single-photon emission has been recently observed around 623 nm from localized defects in both h-BN monolayers and multilayers at room temperature [190]. These polarized emission centers are suggested to serve as room-temperature quantum emitters. The room temperature thermal conductivity of bulk h-BN can reach 400 W m1 K1 while the thermal conductivity of monolayer 2D h-BN is theoretically expected to be more than 600 W m1 K1 [191]. The thermal conductivity of few-layer 2D h-BN is experimentally measured to be 100  270 W m1 K1, making it a promising heat-spreading layer in novel electronic devices or as fillers of polymeric composites. The 2D h-BN can be stable at 1500oC in air and will not react with most chemicals, which makes it one of the thinnest coatings ever shown to withstand extreme environments [192]. Besides strong resistance to oxidation, monolayer h-BN can also serve as a perfect coating to improve the friction performance of the substrates. This high thermal and chemical stability along with excellent coating properties makes h-BN a promising multifunctional coating material.

FIG. 4.28 (A) The left panel is the optical image of a graphene-h-BN lateral heterostructure with electrodes contacting graphene strips (outlined by dotted lines). The right panel shows the I–V characteristics of the indicated devices, with graphene showing conducting behavior and h-BN exhibiting insulating characteristics. (B) The upper left panel is a schematic of a multiple transfer process for ultraflat three-dimensional interconnects. The lower left panel, right panel, and inset are optical images of increasing magnification of a final device substrate with each layer contacted by electrodes. (C) Responsivity of the h-BN detector under 20 V as a function of the incident light wavelength. The inset shows the schematic diagram of the device structure of the h-BN-based deep UV photodetector. (D) Logarithmic scaled I–V curves of the h-BN photodetectors measured in the dark and under 212-nm laser irradiation. (E) Time-dependent photocurrent of the h-BN-based detector at a bias voltage of 5 V, demonstrating good reproducibility and high robustness. (A) and (B) Reproduced with permission from M.P. Levendorf, C.J. Kim, L. Brown, P.Y. Huang, R.W. Havener, D.A. Muller, J. Park, Graphene and boron nitride lateral heterostructures for atomically thin circuitry, Nature 488 (2012) 627. Copyright 2012, Nature Publishing Group. (C) and (D) Reproduced with permission from H. Liu, J.H. Meng, X.W. Zhang, Y.N. Chen, Z.G. Yin, D.G. Wang, Y. Wang, J.B. You, M.L. Gao, P. Jin, Highperformance deep ultraviolet photodetectors based on fewlayer hexagonal boron nitride, Nanoscale 10 (2018) 5559. Copyright 2018, Royal Society of Chemistry. (E) Reproduced with permission from H.L. Wang, X.W. Zhang, H. Liu, Z.G. Yin, J.H. Meng, J. Xia, X.-M. Meng, J.L. Wu, J.B. You, Synthesis of largesized single-crystal hexagonal boron nitride domains on nickel foils by ion beam sputtering deposition, Adv. Mater. 27 (2015) 8109. Copyright 2015, John Wiley and Sons.

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