From h‐BN to Graphene: Characterizations of Hybrid Carbon‐Doped h‐BN for Applications in Electronic and Optoelectronic Devices

Hybrid two‐dimensional materials consisting of graphene and hexagonal boron nitride (h‐BN) have drawn significant interest due to their tunable bandgap and electrical properties. Considering their composition‐dependent properties, ohmic current injection and the development of h‐BN‐based optoelectronic and high‐power electronic devices should be achievable by controlling the C concentration. In this study, electrical and optical characterizations of single‐crystal h‐BN synthesized under high‐pressure and high‐temperature (HPHT) are conducted by varying C concentrations via post‐growth diffusion. Low C‐doped h‐BN (h‐BN:C) with ≈0.1 at% C exhibits nonohmic conduction within a voltage range of ±100 V at all temperatures. In contrast, high h‐BN:C (≈10 at% C) containing C domains and graphite/graphene layers shows additional luminescence peaks and initially exhibits nonohmic conduction at 298 K, which then transforms to ohmic conduction after breakdown‐like behavior at 598 K. This phenomenon, observed only in the high h‐BN:C devices, is attributed to the C‐containing conductive path formed on the channel surface through C drift and local dielectric breakdown of h‐BN mother phase, indicating that ohmic conduction itself does not guarantee the current flow in the conduction/valence bands in h‐BN:C. With these findings, the present thorough and fruitful characterizations are beneficial for the development of h‐BN:C‐based devices.


Introduction
Graphene and h-BN are widely studied 2D materials used in the field of 2D electronics owing to their unique electronic properties. [1,2] Despite the analogous hexagonal structure, graphene is a semimetal with degenerate electronic states near its Fermi energy, [3] while h-BN is known for its insulating nature with a band gap (E g ) of 5.9 eV. [4] Since the first report of the exceptionally high carrier mobility of graphene, [1] numerous attempts have been made to utilize graphene as a semiconductor in field-effect transistors. Nevertheless, a band gap opening is demanding due to its zero-band gap, and one noteworthy method of modification is substitutional doping. The introduction of B and N atoms to engineer the electronic properties of graphene to overcome the drawback of its semimetallic properties has been intensively investigated both experimentally and computationally. [5][6][7][8][9][10][11][12][13][14][15] Combining h-BN with graphene to form an h-BNC film has been calculated to exhibit a nonzero band gap while the carrier mobility can remain high, [5] and band gap tuning can be achieved by varying the concentrations of B and N. [6] Furthermore, the drain current modulation of h-BNC with 40 at% C has been experimentally demonstrated by changing the gate voltage. [8] This study has also shown that such an h-BNC film consisted of segregated BN and C domains with the same hexagonal structure. The hybrid structure formed due to the lowest formation energy of B-N and C-C bonds and a slight lattice mismatch. [8,16,17] The right side of Figure 1 illustrates the reduction in electrical conductivity with the formation of h-BN domains in the direction from graphene to h-BN.
On the other hand, ultrawide band gap h-BN possessing a robust light-matter interaction, [18,19] high dielectric breakdown field, [20][21][22][23][24][25] high thermal conductivity and stability, [26][27][28] and high strength [29] has been extensively utilized as an ideal substrate, [1] a tunneling barrier [30] and a capping layer in 2D electronic devices. Moreover, it also has been demonstrated to be a potential material for deep UV emitters. [18] h-BN is intrinsically an insulator, showing no current under an applied voltage of 100 V, Figure 1. Summary of the atomic structure and electrical conductivity of 2D materials containing B, N, and C, from pure h-BN (leftmost) to pure graphene (rightmost). Below each material shows current and expected applications as an electrically driven device. An irreversible transformation from nonohmic to ohmic conduction at high temperatures is observed only in the high h-BN:C with C domains in the h-BN mother phase.
as depicted on the left side of Figure 1. In recent years, research attention has focused on introducing C into h-BN, moving in the opposite direction from h-BN to graphene, because h-BN:C has been proposed as a promising material for emerging technologies, for example, UV photodetectors/emitters, [31] quantum emissions, [32][33][34] power electronic devices, [35] electrochemical sensors, [36] and electrocatalysts. [37,38] Analogous to those studies on the band gap opening of graphene, a reduction in the E g of h-BN:C as well as its modified electrical and optical properties can be obtained by varying the C concentration, as C creates electronic states in the E g of h-BN. [37][38][39][40] Introduction of a small amount of C may replace B or N depending on the environment. [38] Various methods have been employed to identify the C positions and their energy states. [41][42][43] Nevertheless, a conclusion has not yet been reached. In addition, the states resulting from C substitution have been reported as an origin of photostable and visible single-photon sources at room temperature (RT). [32][33][34]44] Although these optical characteristics of h-BN:C have been extensively studied, emitter applications are desired in the form of electrically driven devices. However, the device development is hindered by the difficulty of carrier injection into high-quality h-BN:C. Nonohmic conduction in the single-crystal h-BN:C under an applied bias of 100 V, as shown in Figure 1, has been achieved in our recent study. [35] Still, ohmic current injection is a major challenge because computational studies have predicted the difficulty of carrier activation due to deep electronic states of ≈1 eV for typical substitutional dopants, including C, in h-BN. [35,45] Instead of C substitution, tuning the electrical conductivity of h-BN by introducing energy states via embedded graphene domains has been predicted in the hybrid h-BNC structure. [6,46] The electronic properties, such as E g , can even be modified by tailoring the graphene structure, suggesting that utilizing the hybrid h-BNC is an effective method to achieve ohmic current injection into h-BN. Nevertheless, modified electrical conductivity of such materials has, thus far, yet to be experimentally demonstrated, and its potential in device applications has not been addressed.
Herein, the electrical conductivity of single-crystal h-BN:C (≈10 at% C) with hybrid structure is investigated from the viewpoint of an application in high-voltage/high-temperature electronic devices. The hybrid h-BNC intrinsically exhibits no current under a small electric field (, verifying its insulating property, and retains the transverse dielectric breakdown field (E BD ) of h-BN. Following contact formation, nonohmic conduction is initially observed at RT. However, at high temperatures, the transformation to ohmic conduction is found, as summarized in Figure 1. The carrier injection mechanism and the origin of transformation to ohmic conduction are investigated to understand the changes in the hybrid h-BNC. Finally, the perspective toward feasible applications is discussed.

Material Characterizations
C doping in h-BN has been generally carried out during the growth of h-BN; [7,31,40,47] nonetheless, the resulting crystallinity is unsatisfactory for device applications. In this study, to maintain high crystal quality, [41] C was introduced into undoped HPHT h-BN crystals after growth. Doping was performed at 3100°C and 2.5 GPa for 30 min, which was more harsh than the previously reported condition at 2000°C and 1 atm for 4 h. [35] This condition resulted in high h-BN:C, while the reported condition yielded low h-BN:C. The detailed doping process is described in the experimental section. It is noteworthy to emphasize that the formation mechanism of h-BN:C differs between localized agglomeration of carbon diffused between layers of h-BN and the separation of carbon once homogeneously distributed as graphite by diffusion through h-BN. The synthesis pathways in the 3100°C region are diverse, and the present results are an evaluation of crystals in which the graphite component is localized in h-BN. Further study is needed to discuss the stability of the solid solution of h-BN and graphite during its formation. The C concentration profiles from secondary ion mass spectrometry (SIMS) are shown in Figure 2a, with a reference from an undoped HPHT h-BN crystal. [35] The high h-BN:C contains ≈10 at% C, which is 100 times more than the low h-BN:C. The optical microscopy (OM) image of the high h-BN:C crystal shows a dark gray color, while the low h-BN:C crystal is yellow. The Raman spectra at 4 K in Figure 2b reveal the apparent difference in bonding characteristics between these two crystals. The E 2g vibrational mode of h-BN appears at 1367.5 ± 0.5 cm −1 for the multilayer flakes of the low h-BN:C, indicating a slight blueshift due to hardening of the force constant at low temperature. In contrast, the E 2g peak of the high h-BN:C varies in the range of 1363-1368 cm −1 depending on the measured position. Apart from the E 2g peak of h-BN, the peaks corresponding to G and 2D modes in graphite/graphene are obviously detected only in the high h-BN:C. The results demonstrate the presence of C-C bonds. Thus, it can be expected that the randomly located C domains are formed in the parent h-BN layer of the high h-BN:C, [8,12,13] as schematically illustrated in Figure 1.
C structures and their distributions inside the high h-BN:C were investigated by cross-sectional transmission electron microscopy (TEM), together with the composition analysis by electron energy loss spectroscopy (EELS), as shown in Figure 2c. The EELS elemental profile along the red line in the TEM image shows the very intense signal of C in the middle and just above the bottom of the flake, while the intensities of B and N at the same positions are weaker compared to the surrounding areas. The EELS mapping suggests that C is present in the form of layers as well as discrete domains in h-BN layers, supported by the EELS mappings of B and N in Figure S1 (Supporting Information). Moreover, the energy loss spectra at those two Cintense positions indicate that the corresponding structure is indeed crystalline graphite due to the clear antibonding * . These results are also consistent with the Raman G peak. However, due to the spatial resolution limit of EELS, the presence of monolayer graphene cannot be identified. Considering the carbon phase diagram under such harsh conditions, [48] it is rational that graphite is present.
In addition, the in-plane distribution of C domains is characterized by conductive atomic force microscopy (c-AFM). Figure 2d presents the topographic and current images of the monolayer high h-BN:C on highly oriented pyrolytic graphite (HOPG). Owing to the fact that only C domains are conductive in the high h-BN:C, the conductive areas in the current image can be regarded as C domains. These results again provide evidence that C domains are separate and randomly distributed in the high h-BN:C. In addition, the sizes of detected conductive domains vary greatly and are typically larger than 20 nm, implying that these domains are probably from the graphite layers observed in the EELS mappings in Figure 2c. Furthermore, the current level of the embedded graphene domains is mainly less than half of the current from HOPG. These results suggest that although the energy loss spectra apparently indicate the crystalline graphite structure, these domains are not completely pure graphene. Instead, small amounts of other elements, for example, B, N, and O, [49] may be included in the domains and possibly are the origin of the unknown peaks in the range of 1700-2000 cm −1 for the Raman spectrum of the high h-BN:C. c-AFM was also employed to observe the monolayer low h-BN:C and undoped h-BN on HOPG. In contrast to the high h-BN:C, conductive domains are not detected in these samples, as shown in Figure S2 (Supporting Information).
The effects of the embedded graphite/graphene domain on the electronic properties of the high h-BN:C are further studied by cathodoluminescence (CL) spectroscopy at room temperature. According to Figure 2e, the CL spectra of the high h-BN:C exhibit additional near UV (NUV) emissions at 387, 412, and 437 nm in addition to the C-related peaks commonly observed in the low h-BN:C. [35,41] The intensities of these peaks are position-and thickness-dependent, similar to those in the Raman spectra of the high h-BN:C. The thick flakes tend to show higher intensity, whereas these peaks are undetectable in the relatively thin flakes ( Figure S3, Supporting Information). Therefore, the additional CL peaks possibly originate from certain structures and impurities, such as the embedded graphite domains and oxygen, in the thick high h-BN:C that do not exist or exist with an insignificant amount in the thin flakes. Interestingly, an early study on the cathodoluminescence of commercial h-BN powders with low purity first reported the similar set of three emission peaks. [50] Recently, the origin for these additional peaks in h-BN powders (98% in purity) annealed in N 2 gas was reported as zero-phonon lines (ZPLs) with their replicas, which are probably attributed to the C-and O-related localized deep-level traps despite their unknown structures. [51] On the other hand, in early studies on single crystal h-BN annealed in a carbon crucible under H 2 gas flow, the effect of H 2 has been discussed. [52,53] Although clear conclusion on the origin of these spectra cannot be drawn at this moment, it is noteworthy that the crystals obtained by diffusing C into the high-quality single-crystal h-BN are the finest material for further investigating emissions.
Here, the peak corresponding to the exciton emissions of h-BN disappears in the high h-BN:C for all flakes. Although the E g reduction is predicted for the hybrid h-BN:C material, [37,38] no evidence is provided from the present CL data. Moreover, it has been predicted that the introduction of the embedded C content can give rise to new states in the E g of h-BN due to the localized p z orbitals of C, [10] which is consistent with the omission of the h-BN exciton peak in the high h-BN:C. Accordingly, the new energy states are proven to be formed in the high h-BN:C. Based on these characterizations, the high h-BN:C is a hybrid material containing a graphite/graphene phase embedded in the h-BN mother phase, as schematically illustrated in Figure 3a.

Device Fabrication and Electrical Characterization
It is necessary to verify the intrinsically insulating nature of the high h-BN:C with graphite/graphene layers. The short-channel devices shown in the inset of Figure 3b are therefore fabricated for the transverse dielectric breakdown experiment. For this purpose, the Ni/Au electrode is deposited without any specific contact formation. It should be noted that quartz is utilized as a substrate due to its E BD (≈6.0 MV cm −1 ) that is greater than that of h-BN and that the electrodes with rounded shapes are prepared to prevent electric field concentration at the corner of the electrode. [21] As presented in Figure 3b, the dielectric breakdown of the high h-BN:C occurs at an applied electric field (F) of ≈0.3 V nm −1 , which is comparable to those of the low h-BN:C [35] and undoped HPHT h-BN. [21] That is, the high h-BN:C is intrinsically insulating, verifying that the graphite layers are discrete and that the h-BN parent phase retains its insulating properties.
Next, the current injection into the high h-BN:C is considered. In our previous study of current injection into the low h-BN:C, [35] introducing defect states in E g and high-temperature treatment initiating metal and h-BN interactions were effective for carrier injection, in addition to band bending at the Schottky junction expected from C doping at ≈10 19 atoms cm −3 . It should be noted that the carrier concentration has not yet been revealed since the Hall measurement is not applicable due to the low conductivity of low h-BN:C. Following such methods, metal contacts are formed with Ar plasma at the contact regions followed by Ni/Au deposition and post-metallization annealing. The OM image in the top-left inset of Figure 4a shows a typical device after contact formation. The channel length is designed to be 3 μm so that F due to the applied voltage of ±100 V is less than the breakdown field of h-BN under the simple assumption that for an insulator, the entire voltage drop occurs in the channel with a uniform F. [54] The temperature-dependent current-voltage (I-V) characteristics of the high h-BN:C in Figure 4a exhibit nonohmic conduction within an applied voltage of ±100 V at all temperatures. The bottom-right corner inset of Figure 4a also reveals that the current flows even at 10 K. This behavior is similar to that of the low h-BN:C case ( Figure S4, Supporting Information). [35] Moreover, the current increases with temperature, implying the importance of the thermal activation process on carrier transport. [55] The thermal activation energies (E a ) extracted from the Arrhenius plot in Figure S5 (Supporting Information), depend on the applied voltage, especially at high temperatures. These results indicate that the dominant energy barrier for carriers to escape from a trapped state can be reduced by the applied voltage, which is in accordance with the Poole-Frenkel (PF) model. [56] The linear relation in the PF plot shown in the upper graph of Figure 4b supports PF conduction in the high h-BN:C device. Because the extracted E a (≈80-130 meV) is much smaller than the Schottky barrier height (SBH, Φ B ≥ 2 eV), the present nonohmic conduction is accordingly attributed to defect-assisted PF conduction [56,57] in the highly defective h-BN at the contact regions, as schematically depicted in Figure 4c. Here, n-type conduction is assumed due to C diffusion in an N-rich environment. [34,38] Interestingly, when the sample temperature is increased and maintained at 598 K, the I-V characteristics of the high h-BN:C transform to ohmic conduction. A large current can be observed within an applied voltage of ±1 V at 598 K, as shown in Figure 4d. The ohmic I-V characteristics remain after the temperature returns to 298 K ( Figure S6, Supporting Information). Since none of the low h-BN:C devices show this transformation during the high-temperature and high-voltage measurements, this transformation is probably related to the randomly distributed C structures inside the high h-BN:C. To investigate the transport mechanism in the high h-BN:C devices after the transformation, the temperature-dependent I-V measurements are performed again at temperatures ranging from 20 to 298 K. According to Figure 4e, the current increases as the temperature increases, and slight nonohmic conduction at low temperatures gradually turns into apparent ohmic conduction at high temperatures. These results suggest that the transport mechanism is still controlled by the thermal activation process. Thus, the Arrhenius plot is employed to extract E a at each specified applied bias, as presented in Figure 4f. In contrast to the E a before the transformation, the obtained E a is independent from the applied voltage in a hightemperature region, and the average E a is only ≈10 meV, which is small compared to the thermal energy at RT. On the other hand, E a at low temperatures is ≈0.3 meV at all voltages. These E a are far less than the value of E a before the transformation. Moreover, the PF plot after the transformation in Figure 4b reveals an almost flat slope similar to that of an ideal ohmic device, indicating that PF conduction is no longer the dominant mechanism for carrier transport. Since the SBH at metal/high h-BN:C is quite large, whereas E a becomes much smaller after the transformation, there should be an additional conduction mechanism apart from the carriers tunneling through defect levels introduced externally in the high h-BN:C.

Conduction Path Formation
To explore what happened during the transformation from nonohmic to ohmic I-V characteristics, current-time (I-t) measurements are conducted under a fixed applied bias of 100 V for 60 s per sweep at 598 K. Figure 5a shows that the current slightly increases continuously and, at a certain point, abruptly reaches the compliance of 20 μA and stays at this level for later sweeps. Such an increment of current is similar to the drastic current enhancement due to dielectric breakdown; however, no explosion or current fluctuation is observed in these devices, as shown in Figure S7 (Supporting Information). In addition, the I-t measurements can be repeatedly performed even after the breakdown-like behavior, and I-V characteristics reveal that the electrical conductivity does not return to nonohmic conduction as time passes. This phenomenon is entirely different from the transverse dielectric breakdown in Figure S7c (Supporting Information), and can be observed only when the current is injected into the high h-BN:C where discrete graphite/graphene domains exist in the h-BN layer at high temperature. Therefore, it is implied that ohmic current injection is related to the local dielectric breakdown of the h-BN parent phase located between graphite/graphene domains. It should be noted that the number of sweeps required to observe this transformation varies from device to device.
The devices with ohmic I-V characteristics are then characterized by AFM to further discuss the local changes in the high h-BN:C channel. The AFM phase images in Figure 5b and topographic images in Figure S8 (Supporting Information), display various arbitrary shapes in the high h-BN:C channel after the transformation. These changes can be detected on the surface of the h-BN:C flake because F concentrates mainly on the surface of devices. [21] An increase in the thickness of the h-BN channel can be detected in most devices; however, some devices, including the  Figure 4. The yellow arrow points at an arbitrary shape in the channel of the high h-BN:C device. S and D stand for the source and drain electrodes, respectively. d) Schematic drawing of the surficial conductive path formed across the channel. e) Cross-sectional TEM images taken at the nonconductive and conductive areas indicated by pink and blue arrows in (b), respectively. f) Cross-sectional TEM images at the source contact. The blue dotted frame emphasizes the position where h-BN bends and detaches from the quartz substrate. The detachment of h-BN is observed only near the metal contact, as indicated by the two blue arrows. g) Surficial elemental analysis by AES, compared with the topographic AFM image. The C-containing path bridges both electrodes and serves as an electrically conductive path. device in Figure 5b, exhibit the contrast only in their phase images without significant topographic changes. Furthermore, the shapes in most devices apparently connect the source and drain electrodes, while separate shapes can also be observed in some other devices. The reasons behind these observations should be related to the originally existing C structures as AFM phase images can approximately estimate the C-concentrated areas, while topographic AFM images cannot indicate the size and position of those areas ( Figure S9, Supporting Information).
The electron beam-induced current (EBIC) mapping technique is then employed to identify the electrical conductivity of the arbitrary shape. In general, the carriers generated due to the band bending at the p-n junction or Schottky junction can be detected, and the defects that act as recombination centers are also detected as dark contrast. [58,59] In the present device, the contrast between conductive (bright) and nonconductive (dark) areas is distinguishable without an applied bias, as presented in Figure 5c. The dark source electrode is grounded, whereas the bright drain electrode is connected to a current amplifier. Although the Schottky junction should exist at the metal/high h-BN:C interface, the carrier generation is negligible, which can be understood by the fact that the edges of the h-BN flake are not identified. Instead, the arbitrary shape observed in the AFM image reveals a similar potential to that of the drain electrode, while the other half shows a potential equivalent to that of the source electrode. In addition, the shape and position of the bright area in the h-BN channel correspond well to the shape in the AFM image. These results indicate that the arbitrary shape is conductive but does not electrically connect the source and drain electrodes unless a bias is applied. It should be mentioned that no Au migration is observed in the channel region. Therefore, it can be inferred that the transformation to ohmic conduction in the high h-BN:C device is probably due to randomly distributed C structures that cause local breakdown of h-BN phases and form an electrically conductive path on the surface across the channel, as schematically drawn in Figure 5d. As for comparison, the EBIC mapping of a h-BN device without the conductive path is shown in Figure S10 (Supporting Information). It is obvious that the position or the edges of h-BN cannot be identified, indicating that h-BN before the transformation is not conductive. Furthermore, to experimentally confirm that the arbitrary conductive shape is responsible for ohmic conduction, Ar/O 2 plasma with a power of 200 W is applied to the channel region for 1 min. The typical thickness removed and damaged is ≈5 nm. [35] After Ar plasma treatment, the I-V measurements reveal a reduced current level in some devices, while other devices show a transformation back to nonohmic conduction ( Figure S11, Supporting Information). Such changes may depend on the original thickness of arbitrary shapes.
The cross-section of the device after the transformation is also characterized by TEM and EELS. TEM images at the channel area with a conductive path and the insulating area in Figure 5e indicate no apparent change in the layered structure and thickness even after the transformation. The results agree well with the AFM results in which no physical topographic changes can be apparently detected. In contrast, the bending of the high h-BN:C flake and the detachment from the quartz substrate near the edge of the electrodes are observed in the TEM images at both contact regions, as presented in Figure 5f and Figure S12 (Supporting Information). This bending is likely to be related to the breakdown of the high h-BN:C in the vicinity of the metal-contact areas because F is locally concentrated and enhanced within a distance of ≈50 nm from the edge of the electrode. [21] It is also worth noting that the bending near the contact is not observed in the TEM images of the low h-BN:C devices. Furthermore, EELS elemental mapping in Figure S13 (Supporting Information), shows C drift along the out-of-plane direction, which is consistent with the electric field direction from the electrode contacting spot. The C-drifted trace is noticeable under the edge of the electrode, and its structure is amorphous ( Figure S14, Supporting Information). In addition, C drift exists at the breakdown position. This trace is, nevertheless, less clear because the C signal intensity at one position is determined not only by the strength of F but also by the concentration of the original C sources. These results thus indicate that F induces not only the breakdown of h-BN but also the C drift.
A surficial elemental analysis was then conducted to further confirm that the conductive path originated from these hybridized C structures. The surface-sensitive Auger electron spectroscopy (AES) reveals that the conductive path mainly contains C (>90%) and evidently connect the electrodes as presented in Figure 5g, despite the distinct disconnected pattern observed in the channel by tapping-mode AFM. In contrast, the distribution of B, N, and O at the conductive path is almost below the detection limit ( Figure S9, Supporting Information). Based on the above observation, it can be expected that during the high-voltage and high-temperature measurements, the intense electric field and the Joule heating effect simultaneously induce the C drift, the local breakdown of the high h-BN:C in the vicinity of contact, and the onset of a surficial current flow path formation on the high h-BN:C channel.

Summary and Outlook
The possibility of utilizing h-BN:C with a hybrid structure as an active material in electronic and optoelectronic applications is discussed. Ohmic current injection is achieved in the singlecrystal high h-BN:C. However, this is not counted as the first and crucial step in developing h-BN electronic devices for the reason that ohmic conduction results from the formation of a conductive path on the channel surface, not from the carrier injection into the conduction/valence bands in h-BN:C. On the other hand, the present study fruitfully reveals the graphite/graphene domains, which give rise to the newly found Raman and CL peaks, in high h-BN:C (≈10 at% C). Since the states created by C in h-BN have been considered an origin of photostable and visible singlephoton sources, further studies on this material are motivated and expected to pave the way for the future development of optoelectronic devices.

Experimental Section
Material Synthesis: Pristine h-BN single crystals were synthesized using a temperature-gradient method under high pressure and high temperature. [4] Carbon was introduced to the pristine h-BN crystals by post-growth diffusion in N 2 atmosphere up to 2100°C. Since h-BN is decomposed above 2100°C at ambient pressure, higher temperature treatment to enhance heavily carbon diffusion was carried out under high pressure by using belt type high pressure apparatus, where temperature above 2300°C was extrapolated by the pre-stablished relationship between the applied electric power and measured temperature with W-Re thermocouples. The pristine h-BN crystals were placed in a graphite capsule with graphite powder, heated at 3100°C under 2.5 GPa for 30 min and then treated in a mixture of sulfuric acid and nitric acid at 350°C for 48 h to remove the graphite component. h-BN:C flakes were mechanically exfoliated from these bulk crystals by the Scotch-tape method and transferred onto a quartz substrate. The substrate was subsequently cleaned by annealing at 350°C for 3 h under an oxygen flow.
Material Characterization: The carbon concentration profiles were obtained by SIMS with an acceleration voltage of 5.0 kV. Raman spectroscopy with an excitation laser of 488 nm was employed to examine the bonding characteristics in h-BN:C. The Raman results were fitted by a single Gauss−Lorentz function, and the Si peak for reference was at 520.7 cm −1 . To investigate the existence of C structures inside the high h-BN:C flakes, EELS elemental mapping and electron energy loss spectra were obtained with an acceleration voltage of 200 kV and a beam diameter of ≈0.1 nm . The distribution of C domains was characterized by conductive AFM with a bias voltage of 3 mV, a loading force of ≈50 nN and a scan rate of 0.5 Hz. The Ti/Ir-coated tip with a radius of 25 ± 10 nm and a spring constant of 2.8 N m −1 was used. The samples for CL measurements were prepared by transferring the high h-BN:C flakes on Si substrate and annealing under O 2 atmosphere at 350°C for 3 h for cleaning purposes. No further annealing was done.