**3. Layer-by-layer graphene growth on β-SiC/Si(001)**

For detailed understanding the mechanisms of the surface transformation and layer-by-layer graphene growth on β-SiC/Si(001) in UHV at high temperatures a series of experimental studies with *in-situ* control of the surface atomic and electronic structure during heating have been conducted [102, 109]. **Figure 7** summarizes the β-SiC(001) surface transformations during annealing in UHV, monitored using *in-situ* core-level PES (**Figure 7(a–g)**) and *ex-situ* LEED and STM (**Figure 7(h–l)**).

#### **Figure 7.**

*(a–g) In-situ core-level PES studies of β-SiC/Si(001) during heating in UHV. (a) Temperature of the sample during the PES measurements. (b–g) Time evolution of the C 1s core-level spectra recorded in snapshot regime during heating. A single spectrum taken in the corresponding temperature interval (shown in panel (a)) is presented. (h–l) Evolution of the SiC(001) surface atomic structure probed by LEED and STM. The 3 × 2, 5 × 2, c(4 × 2), and c(2 × 2) reconstructions are consecutively formed on the SiC(001) surface in the temperature range of 800–1300°C before the graphene overlayer formation. Reproduced from Ref. [109] with permission of Elsevier.*

**141**

*Controllable Synthesis of Few-Layer Graphene on β-SiC(001)*

The first steps toward successful graphene synthesis on β-SiC/Si(001) relate to the removal of the protective silicon oxide layer and the fabrication of a contaminant-free SiC(001)1 × 1 surface structure. This reconstruction can be fabricated after outgassing the sample holder and flash-heating the β-SiC/ Si(001) wafers at 1000–1100°C. Then, the fabrication of a graphene overlayer includes the deposition of several monolayers (MLs) of silicon atoms onto the clean, carbon-rich SiC(001)1 × 1 surface and annealing at gradually increasing temperatures. Depending on the quality of the β-SiC thin film grown on the Si(001) wafer, the deposition and annealing cycles can be repeated until

**Figure 7(a–g)** shows the results of the PES experiments with real-time control during the direct-current sample heating with silicon deposited onto SiC(001)1 × 1 surface structure [109]. During the measurements, a current was applied to heat the sample up to 1350°C (**Figure 7(a)**). The C 1*s* core-level spectra were taken in a snapshot mode during the sample heating with an acquisition time of 1s/spectrum, using a photon energy of 750 eV. Six core-level spectra taken at different stages of the surface graphitization are shown in **Figure 7(b–g)**. Two main C 1*s* peak components can be distinguished in the spectra, which change their relative intensity with increasing temperature. Note that the absolute (but not the relative) binding energies of the individual components in this experiment could be modified by the

At lower temperatures (**Figure 7(b)**), a strong peak corresponding to the bulk carbon atoms dominates in the PES spectra. At temperatures above 1200°C (**Figure 7(d–f )**), an additional component (shifted to higher BE) starts to grow, while the relative intensity of the bulk component decreases. The change of the C 1*s* core-level shape corresponds to the carbonization of the top surface layers at high temperatures. At temperatures close to the silicon melting point (1350°C),

to graphene lattice formation (**Figure 7(g)**). *Ex-situ* LEED measurements proved the existence of a graphene overlayer on the β-SiC/Si(001) wafer used for the PES

**Figure 7(h–l)** shows step-by-step LEED and STM studies of the β-SiC(001) surface atomic structure after heating in UHV at various temperatures. They prove consecutive fabrication of different β-SiC(001) surface reconstructions in accordance with Refs. [110, 111]. The LEED and STM data in **Figure 7(h–l)** were obtained

after consecutive heating of the same β-SiC/Si(001) sample in UHV to 1000, 1150, 1200, 1250, and 1350°C, and cooling to room temperature. After annealing at temperatures of 700–1000°C a uniform, Si-rich SiC(001)3 × 2-reconstructed surface with large (001)-oriented terraces is fabricated (**Figure 7(h)**). Increasing the annealing temperature from 1000 to 1250°C leads to consecutive fabrication of the silicon-terminated 5 × 2 (**Figure 7(i)**), c(4 × 2) (**Figure 7(j)**), 2 × 1, and carbonterminated c(2 × 2) reconstructions (**Figure 7(k)**). According to the LEED and STM studies, the most uniform graphene overlayers on β-SiC(001) can be obtained after flash heating (10–20 s) of the c(2 × 2) reconstruction at 1350°C with post-annealing at 600–700°C, which is similar to the method used for the synthesis of graphene on α-SiC [39, 112, 113]. The LEED pattern shown in **Figure 7(l)** reveals sharp substrate spots and 12 double-split graphene spots related to the formation of the few-layer graphene nanodomain network similar to the one presented in **Figures 1** and **3**. The exact number of the graphene layers synthesized on β-SiC(001) during UHV heating could strongly depend on the vacuum conditions, annealing temperature and duration [70]. To uncover the mechanism of the layer-by-layer graphene growth on the β-SiC/Si(001) substrates and find the way to control the number of synthesized graphene layers and preferential nanodomain boundary direction,

hybridization corresponding

*DOI: http://dx.doi.org/10.5772/intechopen.86162*

sharp 1 × 1 LEED pattern is observed.

voltage applied across the β-SiC/Si(001) wafer.

the carbon-carbon bonds undergo a transition to *sp2*

experiments presented in **Figure 7(b–g)**.

*Silicon Materials*

STM (**Figure 7(h–l)**).

one to several monolayers [50, 82–93, 101].

**3. Layer-by-layer graphene growth on β-SiC/Si(001)**

on millimeter-sized samples. Such nanoribbon systems supported on the Si(001) wafers are very promising because the presence of the self-aligned boundaries can provide a sizeable energy gap in graphene [10]. However, for technological applications it is highly desirable to control the thickness of the graphene overlayer and reduce the number of the preferential NB orientations from two to one. Note that thickness of the few-layer graphene synthesized on β-SiC/Si(001) wafers by different groups, utilizing very similar UHV thermal treatment procedures, varied from

For detailed understanding the mechanisms of the surface transformation and layer-by-layer graphene growth on β-SiC/Si(001) in UHV at high temperatures a series of experimental studies with *in-situ* control of the surface atomic and electronic structure during heating have been conducted [102, 109]. **Figure 7** summarizes the β-SiC(001) surface transformations during annealing in UHV, monitored using *in-situ* core-level PES (**Figure 7(a–g)**) and *ex-situ* LEED and

*(a–g) In-situ core-level PES studies of β-SiC/Si(001) during heating in UHV. (a) Temperature of the sample during the PES measurements. (b–g) Time evolution of the C 1s core-level spectra recorded in snapshot regime during heating. A single spectrum taken in the corresponding temperature interval (shown in panel (a)) is presented. (h–l) Evolution of the SiC(001) surface atomic structure probed by LEED and STM. The 3 × 2, 5 × 2, c(4 × 2), and c(2 × 2) reconstructions are consecutively formed on the SiC(001) surface in the temperature range of 800–1300°C before the graphene overlayer formation. Reproduced from Ref. [109] with* 

**140**

**Figure 7.**

*permission of Elsevier.*

The first steps toward successful graphene synthesis on β-SiC/Si(001) relate to the removal of the protective silicon oxide layer and the fabrication of a contaminant-free SiC(001)1 × 1 surface structure. This reconstruction can be fabricated after outgassing the sample holder and flash-heating the β-SiC/ Si(001) wafers at 1000–1100°C. Then, the fabrication of a graphene overlayer includes the deposition of several monolayers (MLs) of silicon atoms onto the clean, carbon-rich SiC(001)1 × 1 surface and annealing at gradually increasing temperatures. Depending on the quality of the β-SiC thin film grown on the Si(001) wafer, the deposition and annealing cycles can be repeated until sharp 1 × 1 LEED pattern is observed.

**Figure 7(a–g)** shows the results of the PES experiments with real-time control during the direct-current sample heating with silicon deposited onto SiC(001)1 × 1 surface structure [109]. During the measurements, a current was applied to heat the sample up to 1350°C (**Figure 7(a)**). The C 1*s* core-level spectra were taken in a snapshot mode during the sample heating with an acquisition time of 1s/spectrum, using a photon energy of 750 eV. Six core-level spectra taken at different stages of the surface graphitization are shown in **Figure 7(b–g)**. Two main C 1*s* peak components can be distinguished in the spectra, which change their relative intensity with increasing temperature. Note that the absolute (but not the relative) binding energies of the individual components in this experiment could be modified by the voltage applied across the β-SiC/Si(001) wafer.

At lower temperatures (**Figure 7(b)**), a strong peak corresponding to the bulk carbon atoms dominates in the PES spectra. At temperatures above 1200°C (**Figure 7(d–f )**), an additional component (shifted to higher BE) starts to grow, while the relative intensity of the bulk component decreases. The change of the C 1*s* core-level shape corresponds to the carbonization of the top surface layers at high temperatures. At temperatures close to the silicon melting point (1350°C), the carbon-carbon bonds undergo a transition to *sp2* hybridization corresponding to graphene lattice formation (**Figure 7(g)**). *Ex-situ* LEED measurements proved the existence of a graphene overlayer on the β-SiC/Si(001) wafer used for the PES experiments presented in **Figure 7(b–g)**.

**Figure 7(h–l)** shows step-by-step LEED and STM studies of the β-SiC(001) surface atomic structure after heating in UHV at various temperatures. They prove consecutive fabrication of different β-SiC(001) surface reconstructions in accordance with Refs. [110, 111]. The LEED and STM data in **Figure 7(h–l)** were obtained after consecutive heating of the same β-SiC/Si(001) sample in UHV to 1000, 1150, 1200, 1250, and 1350°C, and cooling to room temperature. After annealing at temperatures of 700–1000°C a uniform, Si-rich SiC(001)3 × 2-reconstructed surface with large (001)-oriented terraces is fabricated (**Figure 7(h)**). Increasing the annealing temperature from 1000 to 1250°C leads to consecutive fabrication of the silicon-terminated 5 × 2 (**Figure 7(i)**), c(4 × 2) (**Figure 7(j)**), 2 × 1, and carbonterminated c(2 × 2) reconstructions (**Figure 7(k)**). According to the LEED and STM studies, the most uniform graphene overlayers on β-SiC(001) can be obtained after flash heating (10–20 s) of the c(2 × 2) reconstruction at 1350°C with post-annealing at 600–700°C, which is similar to the method used for the synthesis of graphene on α-SiC [39, 112, 113]. The LEED pattern shown in **Figure 7(l)** reveals sharp substrate spots and 12 double-split graphene spots related to the formation of the few-layer graphene nanodomain network similar to the one presented in **Figures 1** and **3**.

The exact number of the graphene layers synthesized on β-SiC(001) during UHV heating could strongly depend on the vacuum conditions, annealing temperature and duration [70]. To uncover the mechanism of the layer-by-layer graphene growth on the β-SiC/Si(001) substrates and find the way to control the number of synthesized graphene layers and preferential nanodomain boundary direction,

*in-situ* high-resolution core-level and angle-resolved photoelectron spectroscopy, LEEM and μ-LEED studies have been carried out [102].

**Figure 8** shows μ-LEED, LEEM *I-V*, ARPES and micro X-ray photoelectron spectroscopy (μ-XPS) data obtained from the same sample region *in-situ* during the high-temperature surface graphitization in UHV. The μ-LEED pattern and C 1s core level spectra taken in bulk- and surface-sensitive regimes from the β-SiC(001)-c(2 × 2) reconstruction prepared prior to graphene synthesis are shown in **Figure 8(a)**. Then, the temperature of the β-SiC/Si(001) wafer was increased and graphene spots were observed in the μ-LEED patterns. The first graphene monolayer (**Figure 8(b)**) was formed after 10 short flashes at temperatures in the range from 1250 to 1300°C and pressures not exceeding 5 × 10<sup>−</sup><sup>9</sup> mbar. For the fabrication of 2 and 3 ML graphene (**Figure 8(c)** and **(d)**), 50 and 100 flash heating cycles, respectively, were applied at 1300°C. The number of the synthesized graphene layers was defined from the number of minima in the low energy part of the electron reflectivity *I-V* curves presented in **Figure 8** (top). The graphs in **Figure 8** (bottom) depict the evolution of the C 1s spectra acquired in normal emission from a circular sample area (d = 2 μm) at 325, 330, 400, and 450 eV photon energies for a SiC(001)-c(2 × 2) reconstruction **(a)**, mono- **(b)**, bi- **(c)**, and trilayer graphene **(d)**. The selected photon energies correspond to different surface sensitivities of the XPS measurements with the highest sensitivity achieved at 325 and 330 eV. The C 1s spectra were decomposed into individual components corresponding to different carbon atom chemical bonds [102]. The results of the C 1s spectra decomposition are presented in **Figure 8** together with the experimental data (black circles) where the red line is the

#### **Figure 8.**

*In-situ LEEM, μ-LEED, ARPES, and XPS showing control over the SiC(001) surface graphitization during heating in UHV. (a) μ-LEED (top) and μ-XPS data (bottom) obtained from the SiC(001)-c(2 × 2) reconstruction. To illustrate the difference between the c(2 × 2) and graphene spectra, the C 1s spectrum from the 1 ML graphene/SiC(001) system is presented (bottom line). (b–d) μ-LEED, LEEM I-V reflectivity spectra, ARPES constant energy maps, and μ-XPS data obtained from one of APDs of the graphene/SiC(001) system at approximately 1 ML (b), 2 ML (c), and 3 ML (d) coverage. Top row: μ-LEED taken from a circular sample area (d = 0.5 μm) using a 44 eV electron beam (left) and LEEM I-V curves demonstrating one (b), two (c), and three (d) minima (indicated by arrows) corresponding to the number of the synthesized graphene layers (right). Middle row: Photoemission angular distribution maps taken at 0.5 eV (left) and 1.5 eV (right) binding energies, measured using a 47 eV photon energy. Bottom row: C 1s core level spectra (black circles) obtained with 330, 400 and 450 eV photon beams for 1, 2 and 3 ML from a circular sample areas (d = 2 μm) and results of the spectra decomposition. Reproduced from Ref. [102] with permission of ACS.*

**143**

*Controllable Synthesis of Few-Layer Graphene on β-SiC(001)*

graphene peak (Gr), the blue line is the bulk SiC peak, the brown line is the surface β-SiC(001)-c(2 × 2) component, the green dashed line is the background and the cyan line is the envelope. One can note that each spectrum displays only two main components. The Gr peak is shifted by ∼1.65 eV toward higher BE relative to the bulk SiC peak located at 282.9 eV. The intensity of the bulk SiC component decreases both with decreasing photon energy and increasing number of graphene layers. No other components except Gr and SiC were detected in the C 1s spectra, confirming the absence of strong chemical interactions between the graphene overlayer and β-SiC, which would provide additional components with higher BE [103]. The intensities of the individual components in the photoemission spectra shown in **Figure 8** can be used as a reference to distinguish between mono-, bi-, and trilayer graphene on the β-SiC/Si(001) wafers using XPS technique only. Measuring the XPS spectra in the normal emission geometry with 330, 400, and 450 eV photon energies, using the fast dynamic-XPS stations with real-time control of the core level spectra shape [109], one can stop the synthesis procedure when a desirable number of graphene

The *in-situ* ARPES and μ-LEED measurements (**Figure 8**) uncover the origin of the 12 singular spots located between 12 double spots in the LEED pattern taken from the trilayer graphene synthesized on β-SiC(001) (**Figure 2(d)**) and explain the mechanism of the layer-by-layer graphene growth. The singular spots in **Figure 2(d)** are aligned with the SiC substrate spots in contrast with the ±13.5° rotated diffraction patterns corresponding to the rotated graphene nanodomain lattices shown in **Figure 3**. The middle row of images in **Figure 8** shows the ARPES intensity constant-energy maps taken at *E = EF* – 0.5 eV and *E = EF* – 1.5 eV as a function of graphene coverage. The ARPES maps prove the conical shape of the Fermi surface for all preferential graphene nanodomain lattice orientations (two non-rotated lattices and four lattices rotated by ±13.5° relative to the two orthogonal <110> directions) at three graphene coverages studied. Notably, both μ-LEED and ARPES maps measured for the 1 ML graphene/SiC(001) system reveal almost the same intensities of the features corresponding to the non-rotated and ± 13.5° rotated domain lattices. The intensity of the diffraction spots and ARPES features corresponding to the non-rotated lattices is systematically suppressed when graphene coverage increases from 1 ML (**Figure 8(b)**) to 2 ML (**Figure 8(c)**), and then to 3 ML (**Figure 8(d)**). The non-rotated graphene lattice orientations are prevailing only at the beginning of the β-SiC(001) surface graphitization. In contrast, when graphene coverage reaches several monolayers, most of the β-SiC(001) surface is covered by nanodomains with four preferential graphene lattice orientations, rotated ±13.5° relative to the two orthogonal <110> directions (**Figure 2(d)**). **Figure 9** shows **(a)** LEEM and **(b–e)** ARPES data obtained from the 1 ML graphene/β-SiC(001) sample. **Figure 9(b)** and **(c)** shows the constant energy ARPES intensity maps measured from different APDs marked as B and C on panel **(a)**. **Figure 9(d)** and **(e)** shows the dispersions obtained by a cut through the experimental data as indicated by the dashed lines in panels **(b)** and **(c)**, respectively. Eighteen identical linear dispersions are clearly resolved, proving that domains with all six preferential lattice orientations at 1 ML graphene coverage exhibit the same

electronic structure typical of free-standing monolayer graphene.

The prevalence of the μ-LEED and ARPES features associated with the nonrotated lattices at sub-monolayer coverages is a key to understand the mechanism of the graphene growth on the β-SiC/Si(001) wafers. **Figure 10(a–c)** illustrates how the non-rotated graphene domain lattice can match the SiC(001)-c(2 × 2) reconstruction. If the lattice parameters of the c(2 × 2) square unit cell (red square) are doubled, it matches well to a slightly distorted square (green lines) connecting carbon atoms of the graphene lattice, which can be laterally translated to cover

*DOI: http://dx.doi.org/10.5772/intechopen.86162*

layers (1, 2, or 3 ML) is synthesized.

*Silicon Materials*

*in-situ* high-resolution core-level and angle-resolved photoelectron spectroscopy,

**Figure 8** shows μ-LEED, LEEM *I-V*, ARPES and micro X-ray photoelectron spectroscopy (μ-XPS) data obtained from the same sample region *in-situ* during the high-temperature surface graphitization in UHV. The μ-LEED pattern and C 1s core level spectra taken in bulk- and surface-sensitive regimes from the β-SiC(001)-c(2 × 2) reconstruction prepared prior to graphene synthesis are shown in **Figure 8(a)**. Then, the temperature of the β-SiC/Si(001) wafer was increased and graphene spots were observed in the μ-LEED patterns. The first graphene monolayer (**Figure 8(b)**) was formed after 10 short flashes at temperatures in the range from

and 3 ML graphene (**Figure 8(c)** and **(d)**), 50 and 100 flash heating cycles, respectively, were applied at 1300°C. The number of the synthesized graphene layers was defined from the number of minima in the low energy part of the electron reflectivity *I-V* curves presented in **Figure 8** (top). The graphs in **Figure 8** (bottom) depict the evolution of the C 1s spectra acquired in normal emission from a circular sample area (d = 2 μm) at 325, 330, 400, and 450 eV photon energies for a SiC(001)-c(2 × 2) reconstruction **(a)**, mono- **(b)**, bi- **(c)**, and trilayer graphene **(d)**. The selected photon energies correspond to different surface sensitivities of the XPS measurements with the highest sensitivity achieved at 325 and 330 eV. The C 1s spectra were decomposed into individual components corresponding to different carbon atom chemical bonds [102]. The results of the C 1s spectra decomposition are presented in **Figure 8** together with the experimental data (black circles) where the red line is the

*In-situ LEEM, μ-LEED, ARPES, and XPS showing control over the SiC(001) surface graphitization during heating in UHV. (a) μ-LEED (top) and μ-XPS data (bottom) obtained from the SiC(001)-c(2 × 2) reconstruction. To illustrate the difference between the c(2 × 2) and graphene spectra, the C 1s spectrum from the 1 ML graphene/SiC(001) system is presented (bottom line). (b–d) μ-LEED, LEEM I-V reflectivity spectra, ARPES constant energy maps, and μ-XPS data obtained from one of APDs of the graphene/SiC(001) system at approximately 1 ML (b), 2 ML (c), and 3 ML (d) coverage. Top row: μ-LEED taken from a circular sample area (d = 0.5 μm) using a 44 eV electron beam (left) and LEEM I-V curves demonstrating one (b), two (c), and three (d) minima (indicated by arrows) corresponding to the number of the synthesized graphene layers (right). Middle row: Photoemission angular distribution maps taken at 0.5 eV (left) and 1.5 eV (right) binding energies, measured using a 47 eV photon energy. Bottom row: C 1s core level spectra (black circles) obtained with 330, 400 and 450 eV photon beams for 1, 2 and 3 ML from a circular sample areas (d = 2 μm) and results* 

*of the spectra decomposition. Reproduced from Ref. [102] with permission of ACS.*

mbar. For the fabrication of 2

LEEM and μ-LEED studies have been carried out [102].

1250 to 1300°C and pressures not exceeding 5 × 10<sup>−</sup><sup>9</sup>

**142**

**Figure 8.**

graphene peak (Gr), the blue line is the bulk SiC peak, the brown line is the surface β-SiC(001)-c(2 × 2) component, the green dashed line is the background and the cyan line is the envelope. One can note that each spectrum displays only two main components. The Gr peak is shifted by ∼1.65 eV toward higher BE relative to the bulk SiC peak located at 282.9 eV. The intensity of the bulk SiC component decreases both with decreasing photon energy and increasing number of graphene layers. No other components except Gr and SiC were detected in the C 1s spectra, confirming the absence of strong chemical interactions between the graphene overlayer and β-SiC, which would provide additional components with higher BE [103]. The intensities of the individual components in the photoemission spectra shown in **Figure 8** can be used as a reference to distinguish between mono-, bi-, and trilayer graphene on the β-SiC/Si(001) wafers using XPS technique only. Measuring the XPS spectra in the normal emission geometry with 330, 400, and 450 eV photon energies, using the fast dynamic-XPS stations with real-time control of the core level spectra shape [109], one can stop the synthesis procedure when a desirable number of graphene layers (1, 2, or 3 ML) is synthesized.

The *in-situ* ARPES and μ-LEED measurements (**Figure 8**) uncover the origin of the 12 singular spots located between 12 double spots in the LEED pattern taken from the trilayer graphene synthesized on β-SiC(001) (**Figure 2(d)**) and explain the mechanism of the layer-by-layer graphene growth. The singular spots in **Figure 2(d)** are aligned with the SiC substrate spots in contrast with the ±13.5° rotated diffraction patterns corresponding to the rotated graphene nanodomain lattices shown in **Figure 3**. The middle row of images in **Figure 8** shows the ARPES intensity constant-energy maps taken at *E = EF* – 0.5 eV and *E = EF* – 1.5 eV as a function of graphene coverage. The ARPES maps prove the conical shape of the Fermi surface for all preferential graphene nanodomain lattice orientations (two non-rotated lattices and four lattices rotated by ±13.5° relative to the two orthogonal <110> directions) at three graphene coverages studied. Notably, both μ-LEED and ARPES maps measured for the 1 ML graphene/SiC(001) system reveal almost the same intensities of the features corresponding to the non-rotated and ± 13.5° rotated domain lattices. The intensity of the diffraction spots and ARPES features corresponding to the non-rotated lattices is systematically suppressed when graphene coverage increases from 1 ML (**Figure 8(b)**) to 2 ML (**Figure 8(c)**), and then to 3 ML (**Figure 8(d)**). The non-rotated graphene lattice orientations are prevailing only at the beginning of the β-SiC(001) surface graphitization. In contrast, when graphene coverage reaches several monolayers, most of the β-SiC(001) surface is covered by nanodomains with four preferential graphene lattice orientations, rotated ±13.5° relative to the two orthogonal <110> directions (**Figure 2(d)**).

**Figure 9** shows **(a)** LEEM and **(b–e)** ARPES data obtained from the 1 ML graphene/β-SiC(001) sample. **Figure 9(b)** and **(c)** shows the constant energy ARPES intensity maps measured from different APDs marked as B and C on panel **(a)**. **Figure 9(d)** and **(e)** shows the dispersions obtained by a cut through the experimental data as indicated by the dashed lines in panels **(b)** and **(c)**, respectively. Eighteen identical linear dispersions are clearly resolved, proving that domains with all six preferential lattice orientations at 1 ML graphene coverage exhibit the same electronic structure typical of free-standing monolayer graphene.

The prevalence of the μ-LEED and ARPES features associated with the nonrotated lattices at sub-monolayer coverages is a key to understand the mechanism of the graphene growth on the β-SiC/Si(001) wafers. **Figure 10(a–c)** illustrates how the non-rotated graphene domain lattice can match the SiC(001)-c(2 × 2) reconstruction. If the lattice parameters of the c(2 × 2) square unit cell (red square) are doubled, it matches well to a slightly distorted square (green lines) connecting carbon atoms of the graphene lattice, which can be laterally translated to cover

#### **Figure 9.**

*LEEM and ARPES characterization of the nanostructured monolayer graphene. (a) DF-LEEM taken from the 1ML graphene/SiC(001) system. (b) and (c) Corresponding photoemission patterns taken for domains B and C in panel (a) at E – EF = 0.5 eV, from a circular sample area (d = 2 μm, hv = 47 eV). (d) and (e) Dispersion of the Dirac cones obtained by a cut through the data as shown by the dashed lines in patterns (b) and (c). Reproduced from Ref. [102] with permission of ACS.*

the entire c(2 × 2) surface by the graphene overlayer (**Figure 10(a)** and **(c)**). The mismatch of these two quadrilaterals is below 2%, which is likely sufficient to initiate the growth of the non-rotated graphene monolayer on SiC(001)-c(2 × 2). Such a small mismatch cannot be found for other possible surface structures. Therefore, the SiC(001)-c(2 × 2) reconstruction is a necessary step for successful high-temperature graphene synthesis on β-SiC(001). This is very similar to results of the STM studies performed on β-SiC(111) [71], where the transition from a typical (√3 × √3)R30° to an intermediate (3/2 × √3)R30° structure matching the graphene (2 × 2) unit cell was observed before the formation of the honeycomb (1 × 1) overlayer.

**145**

**Figure 10.**

*from Ref. [102] with permission of ACS.*

*Controllable Synthesis of Few-Layer Graphene on β-SiC(001)*

The model in **Figure 10** suggests that carbon dimers of the c(2 × 2) reconstruction (indicated by dotted black oval in **Figure 10(a)**) may be considered the smallest building blocks of the non-rotated graphene lattice, since the distance between carbon atoms in the dimers (1.31 Å) is reasonably close to that of the graphene honeycomb lattice (1.46 Å). In order for graphene growth to begin, extra carbon atoms must be present on the c(2 × 2) surface to provide the substantially higher density of carbon atoms in the graphene lattice. Additional carbon atoms are actually observed during the STM studies as random bright protrusions (**Figure 1**) or linear <110>-directed atomic chains decorating the SiC(001)-c(2 × 2) reconstruction (**Figure 10(e)**). These adatoms form chemical bonds with the dimers of the c(2 × 2) reconstruction at high temperatures and initiate the preferential growth of graphene nanodomains with lattices non-rotated relative to the SiC <110> directions. These domains cannot grow to micrometer-scale due to the presence of linear defects on the SiC(001)-c(2 × 2) surface (**Figure 10(e)**) and the mismatch between the c(2 × 2) and graphene lattices producing strain in the overlayer. However, the reasonably small mismatch of the c(2 × 2) and the graphene lattice (**Figure 10(a)**) leads to the prevalence of the two non-rotated lattice variants in the graphene/SiC(001) system until the first monolayer is complete. The next layers presumably grow on top of the first monolayer starting from the linear defects on the surface (either steps or < 110>-directed linear atomic chains), which is supported by the very fast suppression of the non-rotated domain features in the μ-LEED and ARPES maps with increasing graphene coverage (**Figure 8**). The second and third graphene layers can

*Model of few-layer graphene growth on β-SiC(001). (a) A schematic model showing the non-rotated graphene lattice on top of the SiC(001)-c(2 × 2) reconstruction. Carbon and silicon atoms are shown as gray and yellow spheres, respectively, with c(2 × 2) carbon dimers highlighted by red spheres. The red square indicates the c(2 × 2) unit cell, the green square shows the distorted coincidence quadrilateral resembling a doubled c(2 × 2) unit cell. (b) Quasi-3D view of the SiC(001)-c(2 × 2) reconstruction. (c) A model of the graphene honeycomb lattice with the quadrilaterals showing the surface and overlayer cells. (d) A schematic model of the few-layer graphene growth on SiC(001): At the beginning, domains with a non-rotated honeycomb lattice nucleate in accordance with panel (a), then ±13.5°-rotated lattices start to grow from the linear defects, which become the nanodomain boundaries in the nanostructured graphene overlayer. (e,f) Atomically resolved STM images of the SiC(001)-c(2 × 2) surface (e) and trilayer graphene synthesized on a β-SiC/Si(001) wafer (f). Reproduced* 

*DOI: http://dx.doi.org/10.5772/intechopen.86162*

#### **Figure 10.**

*Silicon Materials*

**144**

**Figure 9.**

(1 × 1) overlayer.

*Reproduced from Ref. [102] with permission of ACS.*

the entire c(2 × 2) surface by the graphene overlayer (**Figure 10(a)** and **(c)**). The mismatch of these two quadrilaterals is below 2%, which is likely sufficient to initiate the growth of the non-rotated graphene monolayer on SiC(001)-c(2 × 2). Such a small mismatch cannot be found for other possible surface structures. Therefore, the SiC(001)-c(2 × 2) reconstruction is a necessary step for successful high-temperature graphene synthesis on β-SiC(001). This is very similar to results of the STM studies performed on β-SiC(111) [71], where the transition from a typical (√3 × √3)R30° to an intermediate (3/2 × √3)R30° structure matching the graphene (2 × 2) unit cell was observed before the formation of the honeycomb

*LEEM and ARPES characterization of the nanostructured monolayer graphene. (a) DF-LEEM taken from the 1ML graphene/SiC(001) system. (b) and (c) Corresponding photoemission patterns taken for domains B and C in panel (a) at E – EF = 0.5 eV, from a circular sample area (d = 2 μm, hv = 47 eV). (d) and (e) Dispersion of the Dirac cones obtained by a cut through the data as shown by the dashed lines in patterns (b) and (c).* 

*Model of few-layer graphene growth on β-SiC(001). (a) A schematic model showing the non-rotated graphene lattice on top of the SiC(001)-c(2 × 2) reconstruction. Carbon and silicon atoms are shown as gray and yellow spheres, respectively, with c(2 × 2) carbon dimers highlighted by red spheres. The red square indicates the c(2 × 2) unit cell, the green square shows the distorted coincidence quadrilateral resembling a doubled c(2 × 2) unit cell. (b) Quasi-3D view of the SiC(001)-c(2 × 2) reconstruction. (c) A model of the graphene honeycomb lattice with the quadrilaterals showing the surface and overlayer cells. (d) A schematic model of the few-layer graphene growth on SiC(001): At the beginning, domains with a non-rotated honeycomb lattice nucleate in accordance with panel (a), then ±13.5°-rotated lattices start to grow from the linear defects, which become the nanodomain boundaries in the nanostructured graphene overlayer. (e,f) Atomically resolved STM images of the SiC(001)-c(2 × 2) surface (e) and trilayer graphene synthesized on a β-SiC/Si(001) wafer (f). Reproduced from Ref. [102] with permission of ACS.*

The model in **Figure 10** suggests that carbon dimers of the c(2 × 2) reconstruction (indicated by dotted black oval in **Figure 10(a)**) may be considered the smallest building blocks of the non-rotated graphene lattice, since the distance between carbon atoms in the dimers (1.31 Å) is reasonably close to that of the graphene honeycomb lattice (1.46 Å). In order for graphene growth to begin, extra carbon atoms must be present on the c(2 × 2) surface to provide the substantially higher density of carbon atoms in the graphene lattice. Additional carbon atoms are actually observed during the STM studies as random bright protrusions (**Figure 1**) or linear <110>-directed atomic chains decorating the SiC(001)-c(2 × 2) reconstruction (**Figure 10(e)**). These adatoms form chemical bonds with the dimers of the c(2 × 2) reconstruction at high temperatures and initiate the preferential growth of graphene nanodomains with lattices non-rotated relative to the SiC <110> directions. These domains cannot grow to micrometer-scale due to the presence of linear defects on the SiC(001)-c(2 × 2) surface (**Figure 10(e)**) and the mismatch between the c(2 × 2) and graphene lattices producing strain in the overlayer. However, the reasonably small mismatch of the c(2 × 2) and the graphene lattice (**Figure 10(a)**) leads to the prevalence of the two non-rotated lattice variants in the graphene/SiC(001) system until the first monolayer is complete. The next layers presumably grow on top of the first monolayer starting from the linear defects on the surface (either steps or < 110>-directed linear atomic chains), which is supported by the very fast suppression of the non-rotated domain features in the μ-LEED and ARPES maps with increasing graphene coverage (**Figure 8**). The second and third graphene layers can

start to grow from the linear defects line-by-line [114] which define the positions and orientations of the nanodomain boundaries in the few-layer graphene/β-SiC(001) (**Figure 10(f)**). In this case, it is energetically favorable for graphene lattices in neighboring nanodomains to be rotated by 27° relative to one another, as the model in **Figure 10(d)** (bottom part) illustrates. The comparison of the atomic resolution STM images of the SiC(001)-c(2 × 2) and trilayer graphene/SiC(001) clearly shows the coincidence of the carbon atomic chain directions in the former structure (**Figure 10(e)**) and nanodomain boundary directions in the latter (**Figure 10(f)**). This result suggests that controlling the density and orientation of defects on β-SiC/Si(001) (e.g., steps on vicinal substrates) could allow the average size of the graphene domains and their orientation to be tuned. This can open a way for synthesis of self-aligned graphene nanoribbons supported by the technologically relevant β-SiC substrate.
