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## Creator

[Ryotaro Sakakibara](https://orcid.org/0000-0001-7150-2831), [Kaito Hirata](https://orcid.org/0000-0003-1489-2938), [Yasufumi Takahashi](https://orcid.org/0000-0003-2834-8300), Wataru Norimatsu, [Yasumitsu Miyata](https://orcid.org/0000-0002-9733-5119)

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This document is the Accepted Manuscript version of a Published Work that appeared in final form in Nano Letters, copyright © 2025 American Chemical Society after peer review and technical editing by the publisher. To access the final edited and published work see https://doi.org/10.1021/acs.nanolett.5c02492.[In Copyright](http://rightsstatements.org/vocab/InC/1.0/)

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[Improved Strain Engineering of Monolayer Transition Metal Dichalcogenides via Van der Waals Epitaxy on Graphene/SiC(0001)](https://mdr.nims.go.jp/datasets/071d4447-b277-4f81-bcfb-dad9ea6ed4d6)

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1  Improved Strain Engineering of Monolayer Transition Metal Dichalcogenides 1 via van der Waals Epitaxy on Graphene/SiC(0001) 2  3 Ryotaro Sakakibara1,2*, Kaito Hirata3,4, Yasufumi Takahashi4,5,6, Wataru Norimatsu7, and Yasumitsu 4 Miyata1,2* 5  6 1 Research Center for Materials Nanoarchitectonics (MANA), National Institute for Materials Science 7 (NIMS), Tsukuba 305-0044, Japan 8 2 Department of Physics, Tokyo Metropolitan University, Hachioji, 192-0397, Japan 9 3 Department of Physical Science and Engineering, Nagoya Institute of Technology, Nagoya 466-8555, 10 Japan 11 4 Department of Electronics, Graduate School of Engineering, Nagoya University, Nagoya 464-8603, 12 Japan 13 5 Research Institute for Quantum and Chemical Innovation, Institutes of Innovation for Future Society, 14 Nagoya University, Nagoya 464-8601, Japan 15 6 WPI Nano Life Science Institute (WPI-NanoLSI), Kanazawa University, Kanazawa 920-1192, Japan 16 7 Faculty of Science and Engineering, Waseda University, Tokyo 169-8555, Japan 17  18 *Correspondence should be addressed to R.S. (SAKAKIBARA.Ryotaro@nims.go.jp) and Y.M. 19 (MIYATA.Yasumitsu@nims.go.jp) 20  21 ABSTRACT 22 Engineering thermal strain is crucial for tuning the properties and functionalities of transition metal 23 dichalcogenides (TMDs). Thermal strain arises from the thermal expansion coefficients (TEC) mismatch 24 between TMDs and substrates, but conventional substrates often induce inhomogeneous broadening in the 25 electronic structure mainly due to surface roughness and charged impurities. Here, we demonstrate 26 uniform thermal strain in monolayer WSe2 via van der Waals epitaxy on graphene/SiC(0001) substrates. 27 Compared to WSe2 grown on graphite, its photoluminescence peaks show a redshift and linewidth 28 narrowing of about 30%. These results suggest that uniform tensile strain is introduced to WSe2 due to the 29 small TEC of SiC, and interfacial graphene suppresses the inhomogeneous broadening. Furthermore, 30 tensile-strained monolayer MoS2 grown on graphene/SiC exhibits enhanced catalytic activity for the 31 hydrogen evolution reaction. Our findings highlight the potential of graphene/SiC substrate as a platform 32 for improved strain engineering in TMDs, enabling future applications in electronics, optoelectronics, and 33 2  electrocatalysis. 34  35 KEYWORDS: transition metal dichalcogenide, graphene/SiC, van der Walls epitaxy, strain engineering, 36 photoluminescence, hydrogen evolution reaction 37  38 MAIN TEXT 39 Two-dimensional transition metal dichalcogenides (TMDs) have attracted significant attention for 40 their physical properties and electronic applications.1-5 A key feature of TMDs is their direct bandgap in 41 monolayer form, which becomes indirect bandgap in multilayers.1, 6 The direct bandgap in monolayers 42 leads to strong photoluminescence (PL),1, 6, 7 making them ideal candidates for optoelectronic applications. 43 Importantly, the PL properties of TMDs can be modulated by applying strain to their lattice. For instance, 44 applying a tensile strain of approximately 2% to a WSe2 crystal through mechanical bending strongly 45 enhances the PL from the A exciton and narrows its linewidth from 42 meV to 24 meV at room 46 temperature.8 These significant changes in PL properties are attributed to strain-induced band structure 47 modifications, which alter exciton-phonon interactions.8-10 In addition, strain can modulate other key 48 properties of TMDs, including bandgap,11, 12 carrier mobility,13, 14 and catalytic reactivity.15, 16 Thus, strain 49 engineering is crucial for advancing the functional development of TMDs. 50 To date, the lattice strain to TMDs has been introduced by mechanical bending or thermal shrinkage 51 mismatch of the substrate. The mechanical bending technique usually requires transferring TMDs onto a 52 flexible substrate.8, 9 This technique is useful for investigating the physical properties of micrometer-sized 53 samples, because of strain tunability. On the other hand, the thermal shrinkage mismatch is particularly 54 important when TMDs are grown directly on substrates. For example, chemical vapor deposition 55 (CVD),17-19 a widely used technique for growing TMDs directly on arbitrary substrates, offers an 56 alternative approach to strain engineering. During CVD growth, the difference in thermal expansion 57 coefficient (TEC) between the TMD and the substrate induces strain in the TMD because both materials 58 contract at different rates during the cooling process. This method facilitates the introduction of built-in 59 thermal strain in large-area TMD films without the need for external mechanical bending. In a notable 60 study, Ahn et al. demonstrated that thermal strain in monolayer WSe2 can be systematically tuned by 61 choosing different substrate species for its growth, which results in varying PL peak positions associated 62 with the bandgap modulation.19 This strain engineering via CVD is critical for optimizing the physical and 63 chemical properties of TMDs. 64 A challenge in CVD-based strain engineering is that substrate surface imperfections including 65 roughness, charged impurities, and dangling bonds typically induce inhomogeneous broadening in the 66 3  electronic structure of TMDs. For example, WS2 monolayers grown on SiO2 and sapphire substrates show 67 relatively broad PL peaks, whereas sharp peaks are observed for these grown on atomically-flat graphite 68 and hBN substrates.17, 18 To suppress the inhomogeneous broadening in CVD-grown strained TMDs, we 69 focused on using graphene as an intermediate layer between TMD and substrate. In this study, we report 70 a remarkably narrow PL linewidth in tensile-strained monolayer WSe2 grown via van der Waals epitaxy 71 on a graphene/SiC(0001) substrate. We found that WSe2 on graphene/SiC exhibits the PL peaks with a 72 redshift and linewidth narrowing of about 30% compared to that grown on conventional bulk graphite. 73 These results suggest that WSe2 is subjected to tensile strain due to the small TEC of SiC, even with 74 graphene present as an interlayer. We also grew tensile-strained monolayer MoS2 on the graphene/SiC 75 substrate, which exhibited remarkable catalytic activity for the hydrogen evolution reaction (HER). 76  77 As illustrated in Figure 1a, we first prepared a homogeneous graphene/SiC substrate by the thermal 78 decomposition method.20-23 Subsequently, WSe2 crystals were grown on the graphene/SiC substrate by 79 salt-assisted CVD method using a two-zone electric furnace.18, 24, 25 For comparison, we also grew WSe2 80 crystals on bare SiC, bulk graphite, amorphous SiO2, and sapphire substrates under the same condition 81 (Figure S1). Figure 1b shows an optical image of triangular shaped WSe2 crystals grown on a graphene/SiC 82 substrate. These crystals tend to align in two opposite directions due to van der Waals (vdW) epitaxial 83 growth on graphene surface.17, 18, 26, 27 Figure 1c illustrates the orientation relationship between WSe2 and 84 the underlying graphene layer. We note that thick multilayers of WSe2 were grown on a bare SiC substrate 85 (Figure S1), likely due to the difference in surface energy.26 Figure 1d displays a cross-sectional high-86 angle annular dark-field scanning transmission electron microscopy (HAADF-STEM) image of the 87 monolayer WSe2 crystal indicated by arrows in Figure 1b (lower magnification images are provided in 88 Figure S2). The electron incidence was parallel to the [11＿00]SiC direction, which is also parallel to the [1189 2＿0]Gra. direction.22 The monolayer WSe2 is uniformly formed on top of the bilayer graphene/SiC. The 90 lattice of WSe2 can be resolved with reduced contrast (inset), indicating an orientation relationship of [1191 2＿ 0]WSe2//[112＿ 0]Gra.//[11＿ 00]SiC. This observation further supports van der Waals epitaxy of WSe2 on 92 graphene/SiC. The energy dispersive X-ray spectroscopy (EDS) line profiles obtained from this cross 93 section also confirm the formation of WSe2 on graphene/SiC. Importantly, the signal for Se was also 94 detected at the graphene/SiC interface across the observation area, with Se distributed two-dimensionally 95 as a bright contrast in the image. This suggests that intercalation of Se into the graphene/SiC interface 96 occurred during the CVD process. Such intercalation phenomenon has been reported for various other 97 elements such as H,28, 29 O,30 S,31 Ga,32 and Pb,33 which terminate interfacial Si atoms by penetrating the 98 graphene/SiC interface. The detailed characterization of this two-dimensional Se will be reported 99 4  elsewhere. 100 We then examined the PL properties of the WSe2 monolayer, as shown in Figure 1e. Interestingly, a 101 very narrow linewidth of approximately 28 meV was observed for WSe2 grown on the graphene/SiC 102 substrate. To the best of our knowledge, PL linewidths as narrow as 28 meV at room temperature have not 103 been previously reported for CVD-grown WSe2.19, 24, 26, 34, 35 The peak center was located at about 1.60 eV. 104 In contrast, monolayer WSe2 grown on graphite exhibited a broader peak with a linewidth of 105 approximately 42 meV, positioned at 1.63 eV. Furthermore, the peak intensity for WSe2 on graphene/SiC 106 was more than twice than that on bulk graphite. These results clearly demonstrate the strain-induced 107 modulation of the electronic structure by the graphene/SiC substrate. 108  109 Figure 1. (a) Schematic illustration on the growth procedure of graphene and WSe2 on a SiC substrate. 110 (b) Optical image of WSe2 crystals grown on graphene/SiC. (c) Crystallographic orientation relationship 111 between WSe2 and graphene, with the orientations of the underlying graphene and SiC indicated by arrows. 112 (d) Cross-sectional HAADF-STEM image of the WSe2/graphene/SiC stack, with overlaid EDS spectra for 113 Si, C, W, and Se. The inset shows an enlarged image taken at lower contrast, where the WSe2 lattice is 114 visible. (e) Typical PL spectra for monolayer WSe2 grown on graphene/SiC (red) and on graphite (gray). 115  116 To investigate the remarkable PL properties of WSe2 grown on graphene/SiC, we conducted PL 117 mapping measurements. Figures 2a, 2b, and 2c show the PL intensity, peak energy, and full width at half 118 maximum (FWHM) maps obtained over a single grain of monolayer WSe2 grown on graphene/SiC, 119 respectively. Additionally, spectra extracted from three different points are presented in Figure 2d. 120 5  Although redshifted peaks with narrow linewidths were observed in most parts of the grain, variations in 121 peak characteristics were evident, as shown in Figure 2d. Specifically, the peak intensity and position 122 depend on the locations. In contrast, monolayer WSe2 grown on graphite exhibited less variation in peak 123 features, as shown in Figures 2e–2h. To further elucidate the PL results, plots of intensity versus energy 124 and FWHM versus energy extracted from the map data are displayed in Figures 2i and 2j, respectively. In 125 Figure 2j, data obtained from monolayer WSe2 grown on SiO2 and sapphire substrates are also plotted (see 126 Figure S3 for their map images and spectra). Figure 2i shows a negative correlation between PL intensity 127 and peak energy. This negative correlation was also reported in a previous study that investigated PL 128 inhomogeneity in WSe2, where the variation is attributed to strain differences in the WSe2 lattice.36 On the 129 other hand, as shown in Figure 2j, the FWHM shows a positive correlation with peak energy, whereby the 130 FWHM decreases as the peak energy decreases. Previous studies demonstrated that when WSe2 is 131 subjected to uniaxial or biaxial tensile strain via mechanical bending, the FWHM decreases, and a redshift 132 occurs as the tensile strain increases, which will be explained later.8, 9, 37 From these observations, we 133 conclude that the WSe2 crystal grown on graphene/SiC has more pronounced built-in tensile strain than 134 that on graphite and sapphire. Specifically, the typical peak position of 1.60 eV observed in WSe2 on 135 graphene/SiC corresponds to a tensile strain of about 0.4%, based on the discussion in Ref. 19. 136 Furthermore, the narrow PL linewidth suggests that, at least within the laser spot size (~1 μm2), the tensile 137 strain is uniformly applied to WSe2 and the inhomogeneous broadening is significantly suppressed, likely 138 due to the atomically-flat and clean surface of graphene.38 A detailed discussion on the strain direction is 139 provided in Supporting Information. 140 The changes in PL spectra can be understood by the strain-induced modification of the band structure 141 of monolayer WSe2. As schematically illustrated in Figure 2k, tensile strain induces the upshift of the Q 142 valley in the conduction band and the decrease of direct band gap at the K-K valley.8, 9, 37 The former 143 suppresses intervalley scattering of the direct K-K exciton, thereby enhancing the radiative transition 144 probability. Consequently, the PL spectra exhibit a redshift, a narrowing of the linewidth, and an increase 145 in peak intensity. 146 One may think that the shift of the E' mode in Raman spectrum can also be utilized to investigate the 147 strain in WSe2. However, the E' mode of WSe2 on graphene/SiC exhibits nearly uniform characteristics 148 across the grain, and its Raman shift is indistinguishable from that of WSe2 on graphite (Figures S4 and 149 S5). Thus, it is difficult to quantify strain using our Raman measurement system, primarily due to the 150 limitation in wavenumber resolution (see Supporting Information for details). In contrast, the PL signal is 151 more sensitive to strain in WSe2 compared to the Raman signal, making it more suitable for strain 152 evaluation. Indeed, in Refs. 19 and 36, the authors assessed the strain in WSe2 based on the shift of the PL 153 6  peak rather than that of the Raman signal.19, 36 154  155 Figure 2. (a) PL intensity, (b) peak energy, and (c) peak FWHM maps of monolayer WSe2/graphene/SiC. 156 (d) PL spectra at three different points indicated by symbols in (a). (e) PL intensity, (f) peak energy, and 157 (g) peak FWHM maps of monolayer WSe2/graphite. (h) PL spectra at three different points indicated by 158 symbols in (e). (i) Plot of PL intensity versus peak energy and (j) plot of peak FWHM versus perk energy, 159 both extracted from the map data. For peak fitting, we basically used a single Voigt function component, 160 except for the sample grown on the sapphire substrate, which required two components. (k) Schematic 161 band diagram of tensile-strained and unstrained monolayer WSe2. Tensile strain results in a narrowing of 162 the band gap at the K-K valley and the reduction in intervalley scattering owing to the upshift of the 163 conduction band at the Q valley. 164  165 To gain insight into the origin of the tensile strain in WSe2 on graphene/SiC, we examined the Raman 166 spectra of graphene. Figure 3a shows typical Raman spectra of graphene obtained before and after the 167 CVD growth of WSe2. For clarity, the contribution of the SiC substrate was subtracted from both spectra 168 7  (the original spectra are shown in Figure S6). In the spectrum before WSe2 growth (gray), the FWHM of 169 the 2D peak was 38 cm-1, which is characteristic of monolayer graphene.23, 39, 40 Also, as indicated by the 170 arrows, several broad peaks originating from the interfacial carbon layer (so-called zero-layer graphene 171 (ZLG)) between graphene and SiC were observed.23, 29, 41 The 2D peak was positioned at 2712 cm-1, blue-172 shifted from that of pristine graphene (~2680 cm-1), suggesting that the graphene layer is compressively 173 strained.42 This compressive strain is commonly explained by the TEC mismatch between graphene and 174 SiC.42 Overall, these observations are consistent with the typical characteristics of monolayer graphene 175 grown directly on SiC.22, 23, 42 After the CVD growth of WSe2, notable changes occurred in the spectrum. 176 The characteristic features of the ZLG diminished, and the FWHM of the 2D peak increased to 65 cm-1, 177 which corresponds to the formation of bilayer graphene.29, 39, 40 This spectral change suggests that 178 intercalation of a foreign element into the graphene/SiC interface occurred during the CVD process, and 179 ZLG was converted to graphene.28-33 In fact, as described above, the intercalation of Se and the presence 180 of bilayer graphene were confirmed by the STEM observation in Figure 1d. Here, the 2D peak remained 181 at 2719 cm-1, indicating that the graphene layers are still compressively strained after the CVD growth, as 182 illustrated in the inset. The persistence of compressive strain suggests a strong interaction between 183 graphene and SiC substrate, even after the Se intercalation. 184 8   185 Figure 3. (a) Raman spectra of graphene before and after the CVD growth of WSe2. For clarity, the 186 contribution of the SiC substrate was subtracted from both spectra. The inset illustration represents the 187 tensile strain in WSe2 and the compressive strain in the graphene layers. (b) Relationship between the PL 188 peak energy of monolayer WSe2 and the TEC difference between WSe2 and the growth substrates. Data 189 extracted from Ref. 19 are also plotted as triangles. For WSe2 grown on graphene/SiC, data point is plotted 190 based on the TEC of SiC. 191  192 9  Based on the above results, we now discuss the origin of the tensile strain in WSe2 grown on 193 graphene/SiC. As mentioned previously, the built-in strain in CVD-grown WSe2 is governed by the TEC 194 of the underlying substrate. Figure 3b shows the relationship between the PL peak energy of WSe2 (vertical 195 axis) and the TEC difference between WSe2 and the substrates (horizontal axis). In addition to the results 196 obtained in our study, results in Ref. 19 are also plotted as triangles. The most significant tensile strain was 197 observed on the SiO2 substrate, as evidenced by the pronounced redshift of the PL peak. This can be 198 attributed to a large TEC difference between WSe2 (~9.5 ppm/K19) and SiO2 (0.55 ppm/K19). On the other 199 hand, since the TEC of sapphire substrate is close to that of WSe2, little strain was introduced and the PL 200 peak energy was close to that of unstrained WSe2 (1.65 eV7, 19). The PL peak energy and the TEC difference 201 have therefore a linear correlation. Here, when the PL peak position on the graphene/SiC substrate is 202 plotted based on the TEC of SiC (~4.5 ppm/K43), it aligns well with this correlation. In other words, the 203 tensile strain in WSe2 grown on graphene/SiC can be explained by the TEC of SiC. Notably, as mentioned 204 above, the underlying graphene layers remain under compressive strain. Therefore, we infer that, after the 205 CVD process, graphene experienced compressive strain due to the contraction of the SiC substrate, while 206 WSe₂ underwent tensile strain as it contracted more, owing to its larger TEC compared to SiC. 207 In the following, we further discuss the narrow PL linewidths in WSe2 grown on the graphene/SiC 208 substrate. As shown in Figure 3b, WSe2 grown on SiO2 experienced strong tensile strain. Given the band 209 modulation induced by this strain, a significant narrowing of the PL linewidth would be expected in this 210 sample. However, the actual PL linewidths were more than 60 meV, as shown in Figure 2j. In contrast, 211 WSe2 grown on graphene/SiC exhibited a much narrower linewidth of about 28 meV, despite experiencing 212 less strain than the sample on SiO2. These results suggest that graphene with the atomically smooth surface 213 and negligible charged impurities effectively suppressed the inhomogeneous broadening of the PL peak.17, 214 18 This highlights that the flatness and cleanliness of the substrate surface are other critical factors in the 215 CVD-based strain engineering. 216 The tensile strain observed in WSe2 on graphene/SiC suggests the presence of an interaction between 217 them. However, little strain was introduced in WSe2 grown on graphite flakes, despite the large TEC 218 difference (Figure 3b). As reported previously,17, 18 this observation indicates that the interaction between 219 WSe2 and graphite is minimal, likely due to the chemical inertness of the highly crystalline graphite surface. 220 One possible origin of the interaction between WSe2 and graphene/SiC is the chemical bond formation 221 between dangling bonds at the WSe2 edges and the underlying graphene or SiC (see Supporting 222 Information for a detailed discussion, including Figures S7 and S8). Further investigation into the origin 223 of the strain will be the subject of future work. 224  225 10  The characteristic thermal strain in TMD grown on the graphene/SiC substrate can potentially affect 226 properties other than PL. We thus investigated the catalytic activity for the HER in monolayer MoS2 grown 227 on graphene/SiC using high-resolution scanning electrochemical cell microscopy (SECCM).44 MoS2 was 228 chosen due to its particularly high catalytic activity among various TMDs.45 We also characterized 229 monolayer MoS2 grown on graphene/SiC and confirmed the introduction of tensile strain (see Figures S9–230 S12). In contrast to WSe2, the PL peak of MoS2 on graphene/SiC exhibited a broadening of linewidth 231 along with a redshift (Figure S11). This opposite trend is due to the band structure of MoS2, where the 232 intervalley scattering becomes more pronounced with increasing tensile strain.9, 37 The observed tensile 233 strain in MoS2 can similarly be explained by the same rationale as for WSe2, where the strain arises from 234 the relatively large TEC of MoS2 (~7.5 ppm/K46) compared to that of SiC (~4.5 ppm/K43).  235 Figures 4a and 4b show typical HER current mapping images for monolayer MoS2 grown on 236 graphene/SiC and graphite, respectively. Notably, MoS2 on graphene/SiC exhibited exceptionally high 237 catalytic activity near the edges of the grain (see also Figure S13). Figure 4c shows current profiles across 238 the MoS2 grains, where the current around the edges of MoS2 on graphene/SiC (red) is an order of 239 magnitude higher than that from the edges of MoS2 on graphite (blue). Also, the broad distribution of 240 active sites near the edges of MoS2/graphene/SiC is evident. Figures 4d and 4e display overpotential and 241 Tafel slope at various positions, estimated from linear sweep voltammetry. The overpotential at 10 mA/cm2 242 and the Tafel slope at the edge of MoS2 on graphene/SiC (red) reach 550 mV and 124 mV/dec, respectively. 243 On the other hand, for the edge of MoS2 on graphite (blue), these values were 570 mV and 153 mV/dec, 244 respectively. The lower overpotential and smaller Tafel slope at the edge of MoS2/graphene/SiC indicate 245 better catalytic performance. 246 It is well known that the edges of MoS2 exhibit high HER activity,44, 45 as confirmed by the current 247 mapping image of MoS2 on graphite (Figure 4b). However, in the MoS2/graphene/SiC sample, active sites 248 are more broadly distributed near the edges (Figure 4a). According to a previous study, in addition to the 249 presence of defects, band modulation induced by tensile strain further enhances the catalytic activity of 250 MoS2.15 Therefore, the built-in tensile strain in MoS2 grown on graphene/SiC may contribute to its 251 enhanced catalytic performance. Another possibility is that the strain facilitated the formation of defects 252 in MoS2, increasing the number of active sites.16 A more detailed investigation into the distribution of 253 defects and strain is necessary to fully elucidate the origin of the high HER activity. Overall, these results 254 further support the versatility of the graphene/SiC substrate for the growth and strain engineering of TMDs. 255 11   256 Figure 4. HER current mapping images of monolayer MoS2 grown on (a) graphene/SiC and (b) graphite, 257 taken at applied potentials of -0.84 V and -0.79 V versus reversible hydrogen electrode (RHE), respectively. 258 (c) Electrochemical current profiles along the white lines in (a) and (b), highlighting the distinct current 259 values for the MoS2/graphene/SiC sample. (d) Overpotential and (e) Tafel slope for graphite (black), basal 260 plane of MoS2 on graphene/SiC (yellow), edge of MoS2 on graphite (blue), and edge of MoS2 on 261 graphene/SiC (red). 262  263 In summary, we have demonstrated the potential of graphene/SiC substrate for improved strain 264 engineering in TMDs. We successfully introduced uniform thermal strain in monolayer WSe2 using the 265 atomically-flat graphene interlayer between WSe2 and SiC. The synthetic WSe2 exhibited a significant 266 redshift in PL peak and a narrow linewidth of about 28 meV, highlighting the ability of graphene to 267 suppress the inhomogeneous broadening caused by the substrate surface. Furthermore, monolayer MoS₂ 268 grown on graphene/SiC showed enhanced HER activity, showcasing the versatility of this substrate for 269 applications beyond optoelectronics. These results underline the importance of substrate engineering in 270 achieving both strain uniformity and improved functional properties in TMDs. Our findings provide 271 valuable insights into the role of substrate surface and TEC mismatch in strain engineering, and show that 272 graphene/SiC can overcome the limitations posed by conventional substrates. Our strategy can be extended 273 to other 2D material/substrate systems, offering a general platform for strain engineering via interface 274 modulation. As a future direction, studying the exciton dynamics of strained WSe2 is of considerable 275 interest, but the rapid non-radiative relaxation induced by graphene limits such investigation. Introducing 276 12  an atomically-thin insulating interlayer, such as hBN, may help overcome this challenge and allow clearer 277 observation of exciton behavior under controlled strain conditions. 278  279 SUPPORTING INFORMATION 280 Experimental details; Supplementary notes on the strain direction in monolayer WSe2, the limitation 281 of wavenumber resolution in Raman spectroscopy, the origin of the interaction between TMDs and 282 graphene/SiC substrate, and the effect of the step-terrace structure and graphene layer on the PL properties 283 of WSe2; Optical images of WSe2 crystals grown on graphene/SiC, graphite, bare SiC, SiO2, and sapphire 284 substrates; Cross-sectional STEM images of WSe2/graphene/SiC as well as in-plane view of the WSe2 285 lattice; PL map and spectra obtained for monolayer WSe2 grown on SiO2 and sapphire substrates; E' 286 Raman map and spectra for monolayer WSe2 grown on graphene/SiC; E' Raman spectra for monolayer 287 WSe2 grown on graphene/SiC, graphite, SiO2, and sapphire substrates; Raw Raman spectra of 288 graphene/SiC; AFM topography images and height profiles for WSe2 grown on graphene/SiC; PL map 289 and graphene 2D FWHM map for a WSe2/graphene/SiC sample; Optical images of MoS2 crystals grown 290 on graphene/SiC and graphite substrates; AFM topography images and height profiles for MoS2 grown on 291 graphene/SiC; PL spectra and strain analysis for monolayer MoS2 grown on graphene/SiC; Raman spectra 292 and strain analysis for monolayer MoS2 grown on graphene/SiC; HER current mapping images for various 293 monolayer MoS2 grains grown on graphene/SiC. 294  295 ACKNOWLEDGMENTS 296 We acknowledge Takahiko Endo and Shunnosuke Murakami for providing the MoS2 samples. This 297 work was supported by the Foundation of Public Interest of Tatematsu, Research Foundation for the 298 Electrotechnology of Chubu, the Japan Science and Technology Agency (JST), the JST ACT-X 299 (JPMJAX23DH), the JST FOREST Program (JPMJFR213X and JPMJFR203K), the JST CREST 300 (JPMJCR23A4), and Kakenhi Grants-in-Aid (JP22KJ1535, JP25K17917, JP21H05232, JP21H05234, 301 JP22H00283, JP22H00280, JP24H00044, JP22H04957, JP25H00835, JP24H01189, JP24K17708, and 302 JP24K01347) from the Japan Society for the Promotion of Science (JSPS). MANA is supported by World 303 Premier International Research Center Initiative (WPI), MEXT, Japan. 304   305 13  REFERENCES 306 (1) Mak, K. F.; Lee, C.; Hone, J.; Shan, J.; Heinz, T. F. Atomically thin MoS2: a new direct-gap 307 semiconductor. Phys. Rev. Lett. 2010, 105, 136805. 308 (2) Radisavljevic, B.; Radenovic, A.; Brivio, J.; Giacometti, V.; Kis, A. Single-layer MoS2 transistors. Nat. 309 Nanotechnol. 2011, 6, 147-150. 310 (3) Lee, C. 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