bk-2016-1246.ch010

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Chapter 10

Tip-Enhanced Raman Scattering of Local Nanostructure on Large Sheet and Microisland Epitaxial Graphene Grown on 4H–SiC (0001) Sanpon Vantasin,1 Shohei Uemura,1 Yoshito Tanaka,1,2 Daichi Doujima,3 Tadaaki Kaneko,3 and Yukihiro Ozaki*,1 1Department of Chemistry, School of Science and Technology, Kwansei Gakuin University, 2-1 Gakuen, Sanda, Hyogo 669-1337, Japan 2Institute of Industrial Science, The University of Tokyo, 4-6-1 Komaba, Meguro-ku, Tokyo 153-8505, Japan 3Department of Physics, School of Science and Technology, Kwansei Gakuin University, 2-1 Gakuen, Sanda, Hyogo 669-1337, Japan *E-mail: [email protected]

This short review article reports our recent tip-enhanced Raman scattering (TERS) spectroscopy studies on nanostructures (e.g. step, ridge, crack and holes) of epitaxial graphene grown on 4H–SiC (0001). Both large-sheet graphene and single crystal graphene microislands are studied. It was demonstrated that TERS can explore individual nanostructures that normal Raman spectroscopy is unable to resolve, due to the limit in spatial resolution from the diffraction limit. We investigated the effects of nanoridge on local strain with TERS, and for the first time, provide direct evidence to the proposed mechanism in previous researches that nanoridges on epitaxial graphene are formed as a relief from compressive strain. TERS study on graphene island reveals that the island which appears highly homogeneous in micro Raman imaging actually contains nanoscale strain variation caused by ridge nanostructures. The characterization of nanoridges on graphene island strengthen the conclusion from the study of ridge on large-sheet graphene, as the regularity and small size of graphene island minimizes the effects caused by interfering factors.

© 2016 American Chemical Society Ozaki et al.; Frontiers of Plasmon Enhanced Spectroscopy Volume 2 ACS Symposium Series; American Chemical Society: Washington, DC, 2016.


Introduction Graphene has been a matter of great interest as a promising material for chemical sensors, biosensors, and electronic devices because of its extremely unique electronic and mechanical properties including but not limited to, very high carrier mobility, unique band gap, quantum hall effect, Kohn anomaly, and negative thermal expansion coefficient (1–9). Although graphene can be synthesized by various method, considerable attention has been paid on graphene grown epitaxially on SiC. This technique relies on the surface decomposition of SiC into silicon vapor at high temperature. The carbon atom left on the surface then forms into graphene. This process involves graphene nucleation and growth, similar to crystal formation process (10, 11). Therefore, with precisely controlled condition, graphene from this method can be grown into large sheet graphene with minimal defects. These feasible properties, together with the electrical insulation of SiC, allow epitaxial graphene to have a bright potential in electronic applications (12–14). Of note is that graphene grown on either face of SiC marks different properties: graphene grown on Si face (0001 face) of SiC strongly interact to the SiC surface while the effect on the C-face (000 face) is much weaker (15). Grapene grown on C-face has extremely high mobility and unique nanostructures on the sheet, which does not occur on Si-face grown graphene (16). Due to the weaker interaction to SiC, C-face graphene contains more nanostructures such as ridges on its sheet. These nanostructures affect graphene properties such as band gap (17), electron mobility (18), and internal strain (19), and thus important to be studied. There are studies of nanostructures on C-face grown graphene using techniques such as scanning tunneling microscope (STM) (18, 20, 21), atomic force microscopy (AFM) (19), and angle-resolved photoemission spectroscopy (ARPES) (22) to deepen understanding of their origin and properties. However, these techniques do not provide some information that Raman spectroscopy can probe. Raman spectroscopy is a conventional method for graphene characterization. It is very useful to investigate crucial properties of graphene properties, including but not limited to amount of layers, defects, strain, doping, etc. (23–29) However, spatial resolution of conventional Raman spectroscopy cannot exceed the diffraction limit of light at λ/2NA, and therefore not sufficient to probe each nanostructure individually. Raman measurement on a nanostructure would then overwhelm with the signal from neighbour area and thus does not represent the information from the nanostructure. To improve the spatial resolution for Raman spectroscopy, we use TERS to study graphene nanostructures. TERS is a technique which utilizes the enhancement of local electric field from plasmon resonance between laser and metallic nanotip (30, 31). Since the enhanced electric field is confined at the nanoscale tip apex, the Raman signal from this near field enhancement can exceed diffraction limit of light (30, 32, 33). Because of this excellent spatial resolution, TERS has been used to study graphene, especially in the investigation of small features such as local strain (34–36), edge boundaries (24, 37–40), and nanodefects (25, 41, 42). Using TERS, we investigated nanoridges on large sheet and microisland epitaxial graphene grown on C-face of SiC (35, 43). Both studies presented the 228 Ozaki et al.; Frontiers of Plasmon Enhanced Spectroscopy Volume 2 ACS Symposium Series; American Chemical Society: Washington, DC, 2016.


decreased compressive strain on nanoridges compared to the nearby flat areas (strain relaxation). This is the direct evidence to the mechanism purposed by many previous publications, in which nanoridge forms during the cooling down process after the synthesis, due to the difference in thermal expansion coefficient between graphene and SiC (19). This result also provides more understanding to our earlier discovery about the stress/strain variation of graphene in nanoscale (36), by demonstrating a relationship of such strain variation to the graphene topology. The result from graphene microisland is important because its high regularity (definitely single layer, no grain boundary, and obvious growth process) allows it to be a model sample for mechanism discussion. In our graphene island study, the alternative mechanism by Hu et al. (44) in which trapped silicon vapor forms into nanoridge is discussed. Other nanostructures such as nanosteps, sub-micro crack, and sub-micro hole are also investigated by TERS. As with nanoridges, TERS reveal the nanoscale information which normal confocal Raman cannot resolve. These investigations have demonstrated the potential of TERS technique in the characterization of individual local nanostructure on epitaxial graphene.

Experimental Section Epitaxial graphene sample were synthesized by silicon sublimation method from SiC. Briefly, on axis 4H-SiC wafers were put in semi-close TaC container and then annealed at 1800 °C for 15 minutes. Throughout the synthesis the argon atmosphere in the heating chamber was controlled to have 6.67 kPa pressure. This process allows precise control of graphene layer. For graphene island, the SiC wafers were pre-treated by YVO4 laser hole-engraving and Si vapor etching to provide step-free basal plane in the hole. The silicon sublimation process as mentioned above then used to grow graphene with high regularity. The graphene islands nucleated only from the rare defect of SiC substrate in the hole. Therefore, in the hole with only single defect, the graphene islands grown from only one point and thus contain no grain boundary. TERS measurement was done with etch silver tip and 514 nm excitation laser. The tips are controlled by non-contact mode AFM. All TERS measurement was performed with top-illumination top-collection setup with the tip at 45° from sample normal plane. The spatial resolution of far-field (normal confocal) Raman spectrometer is around 500 nm, investigated by Raman mapping of polystyrene nanobeads (not shown). The spatial resolution of TERS is around 50 to 75 nm. More details of the experiment is mentioned in the original manuscripts (35, 43). The similar graphene synthesis process is also elaborately explained in the papers from Kaneko Laboratory (14, 45).

229 Ozaki et al.; Frontiers of Plasmon Enhanced Spectroscopy Volume 2 ACS Symposium Series; American Chemical Society: Washington, DC, 2016.

Results and Discussion TERS Study of Large-Sheet C-Face Epitaxial Graphene


Characteristic of C-Face Epitaxial Graphene A typical TERS spectrum of single-layer epitaxial graphene is shown in Figure 1. The TERS spectra show two most prominent peaks. One is the G band at around 1590 cm-1, which arises from two degenerate phonon, namely iTO phonon and LO phonon around the Γ point in the first Brillouin zone (23). Another prominent peak is the G′ band at around 2700 cm-1 (which is also called as 2D band in many studies). This band is originated from a double scattering process of two iTO phonons around the K point (23). These two bands are resulting from inplane phonons, thus their positions are directly affected by strain in the graphene sheet (28). The spectrum demonstrated characteristic of single layer graphene, for example, > 1 ratio value of G′/G bands, less than 27 cm-1 of G′ band FWHM value, and symmetric shape of G′ band (23). Of note is that the D band (~1350 cm-1), which indicate defects (23), is absent in the TERS spectrum. Therefore, there is extremely low defect in the graphene sheet. Raman peaks from SiC also appear on the TERS spectra. The sharp E2 TO peak at 776 cm-1 and the broad A1 LO peak at 982 cm-1 (C6v symmetry) can be seen clearly. The small multiple peaks at around 1500–1600 cm-1 are from scattering of optical branch phonons (46). The E2 TO peak has a constant position throughout SiC sample. Therefore, it is useful for in a calibration of Raman shift.

Figure 1. Typical TERS spectra of epitaxial graphene presenting spectra measured from tip attached and tip retracted position, and also the subtraction result between the two spectra.



The analysis of AFM images illustrates that except for the nanostructures on the graphene sheet, most area on graphene is extremely flat. The height variation in most area is not more than few nanometers. (Figure 2) Therefore, in the AFM images, the nanostructures clearly stand out from the flat area.

Figure 2. AFM topology image of epitaxial graphene in the area with many nanostructures. Noted that the area outside of nanostructure is very flat.

TERS Spectra of Step Nanostructure Figure 3 present an AFM (Figure 3a) and TERS (Figure 3b, and Figure 3c) investigation of a step nanostructure. This kind of nanostructure is not found in graphene from mechanical cleavage, as it arises from the inherent step structure of the SiC substrate. The step in Figure 3a is ~2 nm high. TERS measurements were performed on five points across the step. All TERS and far-field spectra yield characteristics of single-layer graphene, which is > 1 IG/IG′ ratios and 18 – 23 cm-1 FWHM of G′ band. G′ band from the spectra, shown in Figure 3b and the G′ band position in Figure 3c clearly presents that the step structure does not provide a significant difference on graphene between the both side of flat areas and also on the step. There is also no difference between far-field Raman and TERS spectra.



Figure 3. TERS study on nanostep structure. (a) AFM topology image with line profile of a nanostep structure. (b) G′ band in far-field Raman and TERS spectra indicated by cross marks in (a). (c) G′ position from the Lorentzian fitting of bands in (b). Reproduced from reference (35). Copyright 2014, American Chemical Society



TERS Spectra of Ridge Nanostructure Figure 4 represent result of TERS experiment on ridge nanostructure. Every TERS and far-field Raman spectra indicate a single layer graphene in the same fashion as the previous part. The AFM image of the structure is shown in Figure 4a. Many groups have studied ridge structure on epitaxial graphene using AFM/STM and agreed that their results suggest the ‘strain relaxation mechanism’ for the ridge formation (18–20). The mechanism explained that epitaxial graphene is synthesized at high temperature. Upon cooling down, SiC shrinks faster than graphene due to the mismatch in the thermal expansion coefficient of graphene (-6.0 × 10-6 to -0.8 × 10-6 K-1) (9, 47) and SiC (4.5 × 10-6 K-1) (48, 49). (Actually, the negative coefficient of graphene indicates that graphene expands as the temperature falls, but the ‘shrink faster’ analogy also correct mathematically, and easily helps in the understanding of the process.) Since graphene is pulled compressively as SiC shrink, strong compressive strain is generated in graphene sheet. Once this strain exceed the critical buckling strain (50), graphene form wrinkles ‘nanoridges’ to relax from the compression. The schematic explanation of this process is shown in Figure 5. This mechanism is logical and is well supported by the AFM/STM result of the mentioned manuscript. However, there was still a lack of direct evidence. The true evidence must be the measurement of strain on both the ridge and nearby flat areas, and then compare the difference. Even though Prakash et al. (19) evaluated the strain value on the ridge from the topology, the calculation of strain involves approximations such as circular curve fit of ridge shape and also assumption that ridges deforms from graphene with area equal to the base area of ridge. This evaluation indicates positive strain value, which means (stretching) tensile strain. The direct evidence was presented by our first study (35), which directly measured strain on the ridge by TERS and will be discussed in the next paragraph. The result also indicates a compressive strain, instead of tensile strain, on nanoridge. Figure 4a presents a 6-nm high nanoridge and the TERS measurement points across the ridge. Figure 4b shows G′ bands in far-field Raman and TERS spectra collected from the points. The peak position is displayed in Figure 4c. It is clear that G′ bands in the TERS spectra band show a downshift of by 8.7 cm-1 on the ridge compared to the flat area. On the other hand, in the far-field Raman spectra, the peak positions from ridge and flat area do not significantly experience similar downshift. This is because the signal from the ridge is overwhelmed by the signal from the large neighboring flat area, due to the much larger spatial resolution of far-field Raman. On the nanoridge, TERS and far-field Raman spectra provides G′ band position with 11.1 cm-1 difference in the band position. The relationship of G′ band position and strain have been extensively studied (11, 28, 50, 51). Using the equation from the references, we can quantitatively calculate strain from the band position.



Figure 4. TERS study on graphene nanoridge. (a) AFM topology image of a graphene nanoridge, together with line height profile. (b) G′ band of far-field Raman and TERS spectra from seven measurement points indicated in (a). (c) G′ band position acquired by Loretnzian peak fitting of the bands in (b). (d),(e) Strain values calculated using uniaxial and biaxial model from peak positions in (c), respectively. Reproduced from reference (35). Copyright 2014, American Chemical Society. The strain values in (d) and (e) are changed from the original manuscript due to the change in the calculation. The scale values on the y axis change as a result.



Figure 5. Schematic representation of ‘ridge forming as a strain relaxation’ mechanism. The measured G′ band position in Raman experiment of graphene is affected by strain and doping by the equation (11):

where , , , and are G′ band position from measured spectra, band position from unstrained and undoped graphene reference (2674 cm-1, Berciaud et al.) (52), position shift cause by mechanical stain, and shift caused by carrier doping, respectively. However, the intrinsic doping level in epitaxial graphene is in the range which the effect of doping on G′ band position is negligible (11, 29). Therefore,

Since the term

is defined by:

where are Grüneisen parameter of G′ band in the uniaxial strain model, Grüneisen parameter of G′ band in the biaxial strain model, Poisson’s ratio, and strain value, respectively. Then, we got

, for uniaxial and biaxial strain, respectively. The studies that are the references of the equations have thoroughly discussed the tensile and compressive strain of graphene from various aspects, including the first principle calculations at the atomic scale and the experiments at wafer size scale (28, 53). Therefore, the model should be valid in the scale of nanoridges. It is important to mention that each reference for Grüneisen parameter and Poisson’s ratio does report different values. Elaborately, 3.55 and 2.7 for uniaxial Grüneisen parameter (28, 54). 3.8, 2.7, and 2.8 for biaxial Grüneisen parameter (11, 54, 55). 0.13, 0.186, 0.212, and 0.231 for Poisson’s ratio (28, 53, 56, 57). 235 Ozaki et al.; Frontiers of Plasmon Enhanced Spectroscopy Volume 2 ACS Symposium Series; American Chemical Society: Washington, DC, 2016.


The disagreement in the value of Grüneisen parameter may come from the effect of different substrates used in the references. Two values of Poisson’s ratio (0.13 and 0.186) are of the graphene itself, while another two values are Poisson’s ratio of SiC. We still do not know the reason why each reference reported different Poisson’s ratio of the same material. Our criteria is to choose Grüneisen parameter (54) and from the references with similar system as ours, therefore (11). Poisson’s ratio of SiC (57) is chosen due to the strong attachment between graphene and SiC. The calculated strain value for each point, for uniaxial and biaxial strain model is shown in Figure 4d and Figure 4e, respectively. It can be clearly seen in Figure 4d and Figure 4e that there is negative strain value in every measured point. The negatives values indicate that the strain on every point is compressive. This is common for epitaxial grapheme (53, 58), and goes along well with the explanation of ‘ridge as a strain relaxation’ mechanism mentioned in the introduction part. On the ridge center, it is obvious that compressive strain is significantly weaker (strain value is less negative) compared to flat area (Figure 4d and e). The strain difference between ridge and flat area is 1.6 × 10-3 for uniaxial strain and 6.0 × 10-4 for biaxial strain. This weaker compressive strain on the ridge is, undoubtedly, the direct evidence of ‘ridge as a compressive strain relaxation’ mechanism (18, 19). This was also the first time that strain on an individual nanoridge of epitaxial graphene was quantitatively measured using spectroscopy. In the study of Prakash et al. (19), the calculation of strain using topology from AFM indicates a tensile strain value of 0.044. Our result is important because it is a direct measurement of strain, which suggests that the strain on nanoridge is not tensile, but actually relaxed compressive, compared to neighbor area. A comparison between the strain values at the ridge and flat area to the critical buckling strain is interesting. Critical buckling strain is the maximum compressive strain value that graphene can withstand before a graphene sheet loses its flat shape and gains ripples or curvatures (or in the case of epitaxial graphene in this study, nanoridges). Several studies reported critical buckling strains of around –5.3 × 10-3 to –1.2 × 10-2 (50, 59, 60). The difference in the reported values might come from the difference in sizes and shapes of graphene flake, substrates, or methodologies in the measurement/calculation. The strain value from TERS spectrum at the flat area (almost –5 × 10-3) is much closer to the critical buckling strain values from the references, compared to strain value at the ridge. Noted that there is many factors for the critical buckling strain and the critical value in our sample might be different to the sample in the references. The uniaxial and biaxial strain model is the two extreme cases of strain in only one direction and strain exactly equal in two directions. Schmidt et al. suggested that the definition of uniaxial strain implies two different strain value in perpendicular direction and thus should produces splitting in graphene band (11). Nevertheless, in our case, the small strain would only results in unresolvable splitting (28). Considering the morphology of nanoridges, the strain the direction across the ridge should be partially relieved. Therefore the actual strain in this sample should be somewhere in between both model, instead of purely any model. 236 Ozaki et al.; Frontiers of Plasmon Enhanced Spectroscopy Volume 2 ACS Symposium Series; American Chemical Society: Washington, DC, 2016.


For every measurement point, there is no difference between the position of SiC E2 TO peak in TERS spectrum and far-field spectrum. Therefore, residual compressive strain in graphene has no effect on SiC. A crack sub-microstructure is also investigated (Figure 4 of ref (35)). This crack structure is also a common structure found in epitaxial graphene. (Although occur in only small amount per area in our sample.) The area investigated shows a character of single layer graphene. Since the crack is ~20 nm deep, it is certain that this crack is not only of graphene, but also SiC substrate as well. TERS spectra on the crack center represent much lower intensity ratio between graphene band and SiC band, compared to nearby flat area. This indicates less graphene content in the crack center. This cannot be observed in far-field Raman spectra because the insufficient spatial resolution results in averaging from flat area. With a very similar procedure as crack structure, a hole structure have also been investigated (Figure 6). The AFM image of the structure is presented in Figure 6a. Graphene around this hole is multilayer. In the small area around the hole boundary, TERS spectra show significantly decreased intensity ratio between graphene and SiC bands, while far-field Raman spectroscopy provides almost constant ratio between points (Figure 6b). This demonstrates same thing as the result from sub-micro crack: TERS is a powerful tool for investigation of small structure of graphene which conventional confocal Raman cannot resolve.

Figure 6. (A) AFM topology images and (B) intensity ratio between graphene G band and SiC E2 band from the corresponding points marked in the image. TERS Study of Graphene Microisland Overall Structure of Graphene Microisland Semi-low energy SEM image and AFM image of a graphene microisland is shown in Figure 7. This graphene island is grown on the basal plane in the engraved hole of SiC, as explained in the experimental section. The circular structure in the center of Figure 7a is SiC step terraces and does not represent the number of graphene layer. This SiC structure is very important because such SiC step edge is the nucleation point of epitaxial graphene (10). Since the basal plane preparation is so precise that this step terrace structure is the only structure in the hole, this graphene island is also the only single graphene sheet grown on 237 Ozaki et al.; Frontiers of Plasmon Enhanced Spectroscopy Volume 2 ACS Symposium Series; American Chemical Society: Washington, DC, 2016.


this hole. The growth process of this graphene microisland is therefore obvious: it grew radially from the structure. Since there is no other nucleation point, this graphene island is certainly single crystal (especially for the part outside the structure). As in the case of large graphene sheet dicuessed in the previous suction, ridge nanostructures are also presented on the graphene microisland (Figure 7b and c). Due to the lack of any interfering factor, this graphene island is a perfect model sample for mechanism check.

Figure 7. (a) Semi-low energy SEM image of circular step-terrace structure of SiC, together with a graphene microisland. (b) AFM image on the center of graphene island. (c) Small area scanning AFM image from the dashed square in (b). Reproduced from reference (43). Copyright 2015, PCCP Owner Societies. Equiangular hexagon seems to be the preferred shape of graphene island grown from this process, as all graphene island has such shape. (Figure 8) While the angle of microisland is the same at 120°, the length of each edge seems to be governed by the shape of the step-terrace nucleation site.

Figure 8. (a),(b) Semi-low energy SEM images of other graphene microislands in other holes of the engraved SiC wafer. Reproduced from supporting information of reference (43). Copyright 2015, PCCP Owner Societies. 238 Ozaki et al.; Frontiers of Plasmon Enhanced Spectroscopy Volume 2 ACS Symposium Series; American Chemical Society: Washington, DC, 2016.


Figure 9a, b, c and d illustrate Raman images performed on the graphene microisland. The images are constructed using the intensity of G band, position of G′ band, FWHM of G′ band, and intensity ratio of G/ G′ band, respectively. The homogeneity of G band intensity, narrow G′ band FWHM (26.7 ± 1.0 cm1), and high G′/G intensity ratio (2.83 ± 0.30) on the island confirm the singlelayer property of graphene island (23, 61). The hexagonal shape of microisland graphene is clearly reflected in the G band intensity image.

Figure 9. Raman images of graphene microisland, generated from (A) G band intensity, (B) G′ band position, (C) G´ band FWHM, and (D) the intensity ratio between G and G′ bands. Reproduced from reference (43). Copyright 2015, PCCP Owner Societies.

Tip-Enhanced Raman Spectroscopy of the Graphene Microisland Figure 10 provides the relationships between graphene topology studied by AFM and TERS spectra. G´ band position and FWHM values from the corresponding measurement points in Figure 10 a, b, and c are presented in Figure 10d and e. The difference between nanoridges and flat area in TERS spectra is clear, as the TERS results of flat area are no different from far-field Raman, while the G´ band from every point on nanoridges experience downshift up to 9 cm-1. This shift is very similar to the downshift found on the nanoridge of large 239 Ozaki et al.; Frontiers of Plasmon Enhanced Spectroscopy Volume 2 ACS Symposium Series; American Chemical Society: Washington, DC, 2016.


sheet graphene in previous section. The strain relaxation of the most shifted point can be calculated by using the same procedure as in the previous section to be (1.6 ± 0.5) × 10-3, for uniaxial strain model. The standard deviation of both G´ band positions and FWHMs for TERS on nanoridges is significantly larger than that from flat areas. This indicates that graphene islands which appear strain homogeneous in microscale actually do not in the nanoscale. As in the large sheet graphene, nanoscale strain inhomogeneity comes from the nanoridges. The variation of G´ band positions and FWHM in Figure 10d and e are interesting, as it indicates that each nanoridge has different degree of strain relaxation (and maybe orientation). Perhaps this difference has a relationship with the topology, for example, ridge height or distance from ridge junction. The current data is not sufficient to conclude this as the data points is not enough. Therefore this is left as an open question to be explored.

Figure 10. (a) AFM topology image of a graphene microisland. (b),(c) Smaller area AFM topology image from dash squares with “α” and “β” labels in (a), respectively. (d),(e) G´ band information of the spectra from ridges and flat areas with corresponding numbers in (b) and (c), respectively. Reproduced from reference (43). Copyright 2015, PCCP Owner Societies.

In the previous section, we concluded that the ridge forming mechanism using TERS results of a graphene nanoridge on large-sheet graphene. Ones might doubts the conclusion due to many possible interfering factors. For example, “is it possible that nanoridge is actually a grain boundary?” or “maybe ridge nanostructures affected from underlying SiC structure?”. The result on graphene microisland helps answers all of these questions. The same strain relaxation effect on nanoridge still occurs even though the microisland is free from interfering 240 Ozaki et al.; Frontiers of Plasmon Enhanced Spectroscopy Volume 2 ACS Symposium Series; American Chemical Society: Washington, DC, 2016.


factors such as grain boundary, underlying SiC structure, inherent anisotropic stain of graphene sheet, other long-range effects, etc. The mechanism of strain relaxation is further confirmed. One crucial topic to discuss is the alternative mechanism of nanoridges origin, the accumulation of trapped Si vapor suggested by Hu et al. (44) (Figure 11) In this mechanism, Si vapor from the SiC decomposition is collected under graphene sheet and when enough Si atoms are gathered, they form into nanoridges (Figure 11a). This mechanism sounds possible because it is well known that externally added silicon can diffuse into the graphene-SiC interface and even intercalates the graphene sheet out from SiC substrate (62, 63). However, there are some reasons that this mechanism cannot occur on graphene microisland. The most obvious one is that many nanoridges connect to the edge of the graphene microisland. These ridges are therefore open-end. (Figure 11b) Thus, these ridges should not be able to trap silicon vapor, but the ridges are still formed (and also with compressive strain relaxation). This result suggested that the ‘silicon vapor trapping’ mechanism is not sufficient to explain the ridge formation, while the ‘strain relaxation’ is supported and necessary.

Figure 11. (a) The alternative ‘silicon vapor trapping’ mechanism. (b) Ridges connecting to the edge of graphene island are open-end, therefore unlikely to trap Si vapor.



Conclusions This review article has outlined our two recent studies regarding TERS of SiC grown epitaxial graphene. The first study is concerned with the local nanostructure of epitaxial graphene and the second one is the characterization of epitaxial graphene microisland. In the former TERS and far-field Raman spectra from step, ridge, and crack structures of epitaxial graphene were investigated. The step nanostructure did not show a significant effect on the TERS and far-field Raman spectra. The TERS spectra of the ridge structure yield a lower wavenumber of G′ band by 8.7 cm-1, suggesting a relaxation of compressive strain with respect to the surrounding area. Using G′ band positions in the TERS spectra, the strain difference between the ridge center and flat are was evaluated to be 1.6 × 10-3 and 5.8 × 10-4 for uniaxial and biaxial strain, respectively. This is the direct evidence which confirms the proposed mechanism that nanoridges on the epitaxial graphene results from a relief against compressive strain. In the second study hexagonal epitaxial graphene microislands were investigated by TERS. While the island is homogeneous in the scale of micro-Raman imaging, TERS showed that nanoridges cause a nanoscale variation of strain in both magnitude and direction. By using the G′ band shift of 9 cm-1 in TERS spectra we calculated the strain difference between nanoridge and flat area to be (1.6 ± 0.5) × 10-3. Since the similar strain relaxation are observed for both graphene island and large sheet epitaxial graphene, it is logical that the compressive strain relaxation mechanism is enough for the nanoridges to form, as the Si vapour accumulation mechanism is unlikely to occur for this highly regular sample. These studies have illustrated the strength of TERS in the characterization of small and local nano/sub-micro structures on epitaxial graphene.

Acknowledgments We would like to give our gratitude to Adaptable & Seamless Technology Transfer Program through Target-driven R&D (A-Step) from Japan Science and Technology Agency for a part of funding support in this research.

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