# Unraveling the Heterointegration of 3D Semiconductors on Graphene by Anchor Point Nucleation

Thierno Mamoudou Diallo, Tadeáš Hanuš, Gilles Patriarche, Andreas

Ruediger, Abderraouf Boucherif

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Thierno Mamoudou Diallo, Tadeáš Hanuš, Gilles Patriarche, Andreas Ruediger, Abderraouf Boucherif. Unraveling the Heterointegration of 3D Semiconductors on Graphene by Anchor Point Nucleation. Small, 2024, 20 (15), 10.1002/smll.202306038 . hal-04404236

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# Unraveling the Heterointegration of 3D Semiconductors on Graphene by Anchor Point Nucleation

Thierno Mamoudou Diallo,* Tadeáš Hanuš,* Gilles Patriarche, Andreas Ruediger, and Abderraouf Boucherif*

The heterointegration of graphene with semiconductor materials and the development of graphene-based hybrid functional devices are heavily bound to the control of surface energy. Although remote epitaxy offers one of the most appealing techniques for implementing 3D/2D heterostructures, it is only suitable for polar materials and is hugely dependent on the graphene interface quality. Here, the growth of defect-free single-crystalline germanium (Ge) layers on a graphene-coated Ge substrate is demonstrated by introducing a new approach named anchor point nucleation (APN). This powerful approach based on graphene surface engineering enables the growth of semiconductors on any type of substrate covered by graphene. Through plasma treatment, defects such as dangling bonds and nanoholes, which act as preferential nucleation sites, are introduced in the graphene layer. These experimental data unravel the nature of those defects, their role in nucleation, and the mechanisms governing this technique. Additionally, high-resolution transmission microscopy combined with geometrical phase analysis established that the as-grown layers are perfectly single-crystalline, stress-free, and oriented by the substrate underneath the engineered graphene layer. These findings provide new insights into graphene engineering by plasma and open up a universal pathway for the heterointegration of high-quality 3D semiconductors on graphene for disruptive hybrid devices.

## 1. Introduction

The monolithic heterointegration of 3D bulk materials and 2D layered materials has sparked great interest since their extraordinary intrinsic properties can be conjugated to obtain

unique functionalities arising from the physical stacking of such materials.[1,2] Quasi-van der Waals epitaxy (QVdWE) can theoretically enable epitaxial growth of crystalline 3D semiconductors on 2D materials while circumventing the lattice and thermal mismatch issues, thereby extensively reducing defect density in the epilayers.[3,4] However, while the QVdWE of 2D/3D heterostructures[1] has been well established, the epitaxial growth of single-crystalline 3D on 2D heterostructures is an ongoing challenge.[5-9] This is mainly due to the dissimilar lattice structure of the substrate and epilayer, chemical bonding across the interfaces between 3D and 2D materials,[10] but also due to the low surface energy of 2D materials.[5]

Remote epitaxy (RE) has been proposed as a promising technique to circumvent these limitations and achieve high-quality single-crystalline layers.[11] In this approach, graphene is used as a non-polar interlayer. The epilayer's crystalline orientation is achieved via remote interactions that permeate through graphene. This technique has been applied for the growth of single-crystalline III-V, III-N semiconductors,[11-13] and other materials including metals,[14] complex oxides,[15] Halide perovskites,[16] etc. In recent years, substantial advancements and significant progress in understanding the fundamental physics and principles

T. M. Diallo, T. Hanuš, A. Boucherif
Institut Interdisciplinaire d'Innovation Technologique (31T)
Université de Sherbrooke
3000 Boulevard de l'Université, Sherbrooke, QC J1K 0A5, Canada
E-mail: thierno.mamoudou.diallo@usherbrooke.ca;
tadeas.hanus@usherbrooke.ca; abderraouf.boucherif@usherbrooke.ca

<figure><img src="image_3.png" /></figure>

The ORCID identification number(s) for the author(s) of this article can be found under https://doi.org/10.1002/smll.202306038

© 2023 The Authors. Small published by Wiley-VCH GmbH. This is an open access article under the terms of the Creative Commons Attribution License, which permits use, distribution and reproduction in any medium, provided the original work is properly cited.

DOI: 10.1002/smll.202306038

T. M. Diallo, T. Hanuš, A. Boucherif
Laboratoire Nanotechnologies Nanosystèmes (LN2)–CNRS IRL-3463
Institut Interdisciplinaire d'Innovation Technologique (31T)
Université de Sherbrooke
3000 Boulevard Université, Sherbrooke, Québec J1K0A5, Canada

G. Patriarche
Centre de Nanosciences et de Nanotechnologies (C2N)
CNRS
Université Paris-Saclay
Palaiseau 91120, France

A. Ruediger
Nanoelectronics-Nanophotonics INRS-EMT
1650, Boulevard Lionel-Boulet, Varennes, QC J3×1P7, Canada
<figure><img src="image_4.png" /><figcaption>**Figure 1.** Schematic illustration of Anchor Point Nucleation approach enabling the growth of monocrystalline layers on top of single-layer graphene (SLG). a) SLG is transferred on the semiconductor substrate, b) controlled introduction of defects in SLG by plasma treatment c) Preferential nucleation on SLG d) epitaxial growth of monocrystalline epilayers.</figcaption></figure>

governing RE,[12,16,17] have solidified RE as a prospective method for the fabrication of freestanding membranes. Furthermore, recent studies have suggested that the presence of native oxides beneath the graphene can induce the formation of pinholes within the graphene interface, helping the orientation of the epilayer,[18] in addition to remote polar interactions through the graphene. The quality of the graphene interface is a critical factor influencing the quality of the final membranes. The presence of residues, multilayers, wrinkles, or grain boundaries within the graphene can lead to the formation of polycrystalline layers.[19] To address these issues, several innovative solutions have been recently introduced, such as dry transfer techniques,[19] or the growth of high-quality graphene directly on semiconductor substrates.[20,21] Nonetheless, a significant challenge persists in the realm of RE. Given that the success of RE growth depends on the polarity of both the underlying substrate and the epilayer, this method is exclusively suitable for polar materials.[12] Consequently, it is impractical for applications involving non-polar materials like silicon (Si) and germanium (Ge). Hence, an alternative approach is imperative to alleviate these constraints and enable the growth of non-polar, single-crystalline layers while facilitating the production of freestanding membranes through 2D material-assisted epitaxy.

In this work, we propose a universal approach called anchor point nucleation (APN), based on the controlled introduction of nanoholes in graphene interface by plasma treatment and epitaxial overgrowth, to monolithically grow single-crystalline semiconductors on graphene-terminated substrates (including non-polar materials). We demonstrated that by using plasma treatment, we can introduce defects in the graphene layer that improve its surface reactivity and adhesion with other materials.

Those defects act as preferential nucleation sites for the epitaxial growth as clear evidence was provided by in-situ transmission electron microscopy (TEM) studies from our previous work.[9] Defect engineering using other methods has been previously reported for the epitaxial growth of 2D heterostructures.[22,23] However, this is the first time that such an approach is used for the growth of semiconductors on non-polar materials such as Ge. Moreover, the nature of the induced defects, their role in the epitaxial growth, and the underlying mechanisms for such a technique are elucidated. Here, we find that the induced defects are mainly dangling bonds and nanoholes depending on the treatment duration. Our experimental data provide a clear corroboration for the role played by the induced defect during the nucleation process. We unveil that the APN approach is governed by the nucleation in the nanoholes followed by an overgrowth of a continuous layer. The combination of HRTEM measurements and strain mapping demonstrates that the APN method enables the epitaxial growth of single-crystalline Ge layers without any deformation or defects, caused by lateral overgrowth on graphene-terminated Ge substrate. Our findings demonstrate that the APN approach is a powerful technique that is suitable for the epitaxial growth of 3D semiconductors on graphene and is a promising solution to circumvent the limiting factors of QVdWE and RE.

## 2. Results and Discussion

### 2.1. Anchor Point Nucleation Approach

**Figure 1** shows the schematic representation of the anchor point nucleation (APN) processing steps that enhance the reactivity of graphene and enable the growth of monocrystalline layers
<figure><img src="image_5.png" /><figcaption>**Figure 2.** Analysis of the influence of plasma treatment duration during graphene surface engineering. a) Raman spectra of single-layer graphene for different durations, b) $I_p/I_G$ ratio extracted from the Raman measurements in (a), c) High-resolution XPS spectrum of C1s of a non-treated SLG showing the sp² bonding of the graphene. d) C=C sp² bond area (blue) and sp³/sp² bond ratio (red) extracted from the XPS measurements of C1s orbital for the different treatments.</figcaption></figure>

on single-layer graphene (SLG). This approach is based on the controlled introduction of defects in the SLG, which will act as preferential nucleation sites allowing to anchor adatoms on the graphene surface. The introduced nanoholes then serve as seedling sites to provide the crystal orientation to the epilayer, which is obtained by epitaxial overgrowth on SLG.[18,24] To achieve this, first, large-area SLG is transferred on the substrate using the wet transfer technique (Figure 1a, see Experimental Section for more details). The substrate is then treated by plasma to introduce defects such as dangling bonds and nanoholes in the SLG (Figure 1b). These defects significantly enhance the surface energy of the graphene[25] and serve as preferential nucleation sites for epitaxial growth. Moreover, the nanoholes provide a direct link between the substrate and epilayer enabling substrate-oriented growth even for elemental materials. It is worth mentioning that the use of plasma treatment is more advantageous compared to other patterning techniques such as photolithography and electrolithography,[24,26] which are complexes, expensive, and uses photoresist. In contrast, the plasma-based technique is easy to use, cost-effective, and simple without multistep processes and it enhances the cleaning of PMMA residues from wet transferred graphene,[27-29] preventing random nucleation that will result in 3D growth. After the formation of dangling bonds and nanoholes by plasma treatment, we proceed with the nucle-

ation of semiconductor islands that grow preferentially on the induced defects of the engineered graphene layer (Figure 1c). The nucleation on nanoholes allows the islands to be oriented by the underlying substrate. The nanohole-mediated nucleation is then used as a seed layer to grow a complete epitaxial layer through the lateral overgrowth of the islands (Figure 1d). Compared to QVdWE and RE, the APN approach enables obtention of substrate-oriented epilayers even for non-polar materials (e.g., Si and Ge), maintaining a significant portion of the weak bonding to the substrate, due to the epitaxial overgrowth on graphene interface. This is especially interesting for further exfoliation of the film.

## 2.2. Effect of Plasma Treatment on SLG

A key challenge of the APN approach is to prevent a complete etching of the atomically thin SLG and introduce defects in a controlled manner. To this end, a balance between the duration and the power of the plasma must be established to slightly damage the SLG, while maintaining its nature and coverage on the substrate.[24] This is not an easy task since the graphene can be easily etched away by standard plasma etching techniques.[30] Indeed, our initial studies indicate that a very light plasma
treatment in a short burst must be used to avoid complete re-
moval of the SLG on the surface (details can be found in the Ex-
perimental Section). To further elaborate on these findings the
effects of O₂ plasma etching (320 mTorr O₂ pressure, Radio fre-
quency power of 10 W, and treatment duration varying between
0 and 30 s) on the SLG quality were thoroughly investigated by
Raman spectroscopy and X-ray photoelectron spectroscopy (XPS)
measurements.

**Figure 2a** shows Raman spectra of transferred SLG on the Ge substrate for different durations of plasma treatment. The Raman spectra consist of a set of distinct peaks. The G band is related to C—C bond stretching in sp² carbon systems, the D band is related to defects and disorder, and the 2D band is the second order of the band D, but does not require a defect.[31] The non-treated SLG shows very sharp G and 2D bands with a very low D band, demonstrating the high-quality graphene layer with large crystallite sizes. Let us remember that the ratio of the D to G intensities (I<sub>D</sub>/I<sub>G</sub>) are inversely proportional to the crystallite size.[32] The apparition of D band and significant reduction of 2D band after the application of plasma treatment shows clear evidence of defect introduction inside the graphene layer. The I<sub>D</sub>/I<sub>G</sub> ratio as a function of treatment duration extracted from Figure 2a is shown in Figure 2b. A sharp initial rise of the D band is observed up to 6 s and then it decreases with increasing treatment duration. During the initial stages of the treatment, we believe that mainly dangling bonds and atomic vacancies are introduced in the SLG. As the duration increases larger nanoholes are formed in the SLG causing a significant reduction in the graphene crystallite size. From Figure 2a,b, we concluded that up to 18 s of treatment, the SLG is being transformed from pristine graphene to nanocrystalline graphene with smaller crystallite sizes separated by small amorphous patches.[33,34] The annealing of the substrate during the epitaxial growth transforms the amorphous patches and nanosheets into crystalline graphene as demonstrated by Turchanin et al.[35] Such transformation is characterized by the appearance of D band, the increase of I<sub>D</sub>/I<sub>G</sub> and the broadening of all the peaks[34] (see Figure S1, Supporting Information). Above 18 s of plasma treatment, the nanocrystalline graphene is transformed to low sp³ amorphous carbon.[33]

To further investigate the defect introduced in SLG, we per-
formed XPS to measure spectra of C1s core level and to study
the nature of chemical bonds in engineered graphene. Figure 2c
shows the high-resolution C1s spectrum of a non-treated SLG.
This spectrum contains mainly the sp² carbon peak centered at
284.4 eV. The analysis of the XPS spectra after the plasma treat-
ment is presented in Figure 2d showing the C=C sp² bond area
(blue) and the sp³/sp² bond ratio (red) for different treatment
durations. After, the plasma treatment, the FWHM of the C1s
peak became broader (Figure S2, Supporting Information). From
Figure 2d, we can see that the C=C sp² area decreases with the in-
crease of treatment duration, while the ratio sp³/sp² increases as
more defects are generated in the SLG. This clearly demonstrates
the introduction of defects in graphene. Initially, the sp² area de-
creases rapidly due to the generation of dangling bonds. As the
nanoholes start to grow by merging multiple atomic vacancies,
the slope of the curve of sp² area decreases. These observations
are in good agreement with the Raman spectroscopy measure-
ment presented in Figure 2a,b. More XPS analysis is presented
in Figure S2 and Table S1 (Supporting Information).

2.3. Preferential Nucleation Induced by Defects

To bring new insight into the influence of defects in SLG on epi-
taxial growth and to corroborate our above-mentioned hypothe-
sis, we studied the impact of plasma treatment on epitaxial nuclea-
tion of a few nm of Ge on the treated SLG. On pristine graphene,
the formation of seeds is difficult due to the SLG's low surface en-
ergy, which makes epitaxial nucleation very challenging, because
the adatoms barely stick to the surface and the nuclei can be eas-
ily desorbed before reaching the critical size.[⁹] Such nucleation
results in a polycrystalline 3D growth mode as schematically il-
lustrated by **Figure 3a** (see also Figure S3a, Supporting Informa-
tion). This is further confirmed by grazing incident X-ray diffrac-
tion (GIXRD) measurements (Gray data set). As stated in the pre-
vious part, a very short duration of plasma treatment introduces
mainly dangling bonds and atomic vacancies, which help to in-
crease the wettability of graphene and therefore significantly en-
hance its reactivity. This results in the increase of the Ge (111) ori-
entation, which is the most favorable in the case of QVdW growth
of Ge on SLG,[⁹] as shown by grazing incident X-ray (GIXRD)
diffraction results shown in Figure S4 (Supporting Information).
Even though those dangling bonds enhance the nucleation and
crystal orientation of the epilayer on the SLG by acting as prefer-
ential nucleation sites as represented in Figure 3b, the resulting
layer cannot be epitaxially oriented by the substrate underneath
the graphene layer since there is no direct link between them.
For treatment of 12 s and above, we observe the growth of seeds
with crystal form, which is clear evidence of the nanohole forma-
tion in the SLG (Figure S3c, Supporting Information). In con-
trast to dangling bonds, nanoholes provide a direct link with the
substrate underneath the SLG, enabling direct epitaxial growth
of substrate-oriented seeds (see Figure 3c). Similar observations
have been reported[¹⁸] in nanoholes occurring during the under-
lying substrate deoxidation. The nanoholes-mediated nucleation
must be dominant to grow epitaxially monocrystalline layers on
top of engineered graphene.

By increasing the duration of plasma treatment, more
nanoholes are introduced in the SLG structure, and their size
increases. Figure 3d–f shows SEM micrographs tracking the evo-
lution of the nucleation of Ge on treated SLG at durations of 18,
24, and 30 s, respectively. From these images, we can see that the
density and the size of the seeds are increasing with treatment
duration. The increase in the crystal size is attributed to the pres-
ence of larger nanoholes. For further understanding, we plotted
the surface coverage by the seeds as a function of treatment du-
ration shown in Figure 3g. We found that the surface coverage
is proportional to the treatment duration. This can be explained
by the fact that as the duration increases, larger regions not cov-
ered by graphene become available, enabling the nucleation of
larger seeds directly on the substrate. When the SLG is totally
etched away (above 30 s of treatment), the seeds form a con-
tinuous and homogeneous layer. To successfully perform APN
growth, the presence of the SLG is necessary, as it provides a sur-
face with weak VdW bonding that facilitates the detachment of
the epilayer. For this reason, the plasma treatment duration needs
to remain <30 s. It is also important to note that increasing the
number of direct links between the epilayer and the substrate will
increase the adhesion force of the film and make its exfoliation
more difficult.[24]
<figure><img src="image_6.png" /><figcaption>**Figure 3.** Study of the effect of plasma treatment duration on the nucleation of Ge by APN on engineered SLG on Ge substrate. a–c) schematic illustration showing the different growth mechanisms on pristine graphene, induced by dangling bonds (<12 s of the plasma treatment) and via nanoholes (>12 s of plasma treatment), respectively. On the pristine graphene, a 3D growth mode is observed, while through the nanoholes a single-crystalline layer oriented by the underlying substrate is obtained. d–f) SEM images tracking the evolution of the nucleation of Ge on treated SLG with a duration of 18, 24, and 30 s, respectively. For all the treatments presented in this figure, the growth temperature was 400 °C and the thickness was 5 nm. For short durations (up to 6 s), the APN proceeds through dangling bonds and for longer durations through the nanoholes induced in the graphene layer. g) evolution of the graphene surface coverage by Ge nuclei as a function of treatment duration extracted from the SEM images.</figcaption></figure>

## 2.4. Growth of Ge on Engineered Graphene Using APN Approach

After unraveling the different mechanisms of nucleation involved in the APN approach, we performed the epitaxial growth of Ge layers on the engineered SLG (18 s of O₂ plasma treatment) supported by a Ge (100) substrate. Figure 4a,b, shows plan-view scanning electron micrographs (SEM) of the as-grown epilayer. The SEM images reveal a homogeneous and smooth surface of the 500 nm-thick Ge layers, confirming their high morphological quality. In Figure 4a, we can notice a few defects, which we attributed to the poly-nucleation induced by the multilayer graphene. This can also be seen in Figure 2 showing the SEM

images of the early stages of the nucleation. However, without the multilayer of graphene the epilayer surface is quite homogeneous (Figure 4b). The morphological quality of the epilayers was also investigated by AFM and the resulting image is displayed in Figure 4c. From this figure, we can notice a featureless and smooth surface with an RMS roughness ≈2 nm. The electron backscattering diffraction (EBSD) map of the top surface presented in Figure 4d reveals that the Ge epilayers grown on the engineered SLG are perfectly single-crystalline and have the same crystal orientation as the Ge (100) substrate. This demonstrates that the epilayer is oriented by the underneath substrate through the induced nanoholes. On pristine SLG, the as-grown layers are
<figure><img src="image_7.png" /><figcaption>**Figure 4.** Growth of Ge layer (500 nm) on plasmatreated (18 s) single-layer graphene on Ge (100) substrate using the APN approach. a) Low magnification plan-view SEM image showing homogenous surface with some defects induced by the nucleation on multilayer graphene, b) SEM image of the red square on (a) showing a surface without any defect, c) AFM image of a scan 5 × 5 µm² area showing smooth surface with an RMS roughness of 2 nm, d) EBSD map of the layer showing single orientation, which demonstrates its single-crystalline nature.</figcaption></figure>

polycrystalline as previously reported in the literature,[12,19,36,37] a result of spherical polycrystalline nucleation (Figure S5, Supporting Information). In contrast, the APN approach, which is based on the controlled introduction of nanoholes in SLG and epitaxial overgrowth, has successfully demonstrated the growth of continuous and smooth layers on a non-polar substrate. So far, this is unachievable with the QVdWE or RE[11] since they are governed by the interactions with 2D interface or the substrate underneath the graphene.[12]

To further investigate the crystalline quality of the as-grown layers by the APN approach, cross-sectional TEM measurements were performed. Figure 5 shows the TEM analysis of Ge epilayers grown on engineered SLG. The cross-sectional high-angle annular dark-field STEM (HAADF-STEM) of the epitaxial structure is presented in Figure 5a. From this figure, we can see that the epi structure is made of a continuous Ge epilayer, the engineered SLG interface, and the Ge substrate. We can also notice some small holes (Figure S6, Supporting Information) highlighted by the red arrow, which we attributed to the overgrowth of Ge on the SLG once the nucleation is initiated within the nanoholes. This observation is in good agreement with our proposed mechanism for the APN approach presented in the previous sections. Figure 5b shows the cross-sectional high-resolution TEM image of the epitaxial structure where the SLG interface (dark contrast) can be observed clearly between the epilayer and the substrate. This HRTEM image was obtained with the zone axis <001>. Note

that for a better view, the image was rotated by 90° with respect to the initial image. Through the dark interface representing the perfect graphene layer in Figure 5b, we can notice some small discontinuities (highlighted by yellow arrows) indicating the induced nanoholes. This observation confirms that the SLG still exists after the Ge growth, indicating that the graphene is not damaged under the used growth conditions and still maintains its nature. From this HRTEM image, there is clear evidence of the absence of structural defects in the epilayer, which demonstrates the high structural quality of the as-grown layers by APN. Figure 5c,d, displays the fast Fourier transform (FFT) patterns of the epilayer and the underlying substrate shown in Figure 5b, respectively. These FFT patterns display the family planes of (220) and (400) of the Ge diamond cubic lattice with d-spacing of ≈0.2 and 0.14 nm, respectively. We only observed those family planes due to the fact that the zone axis is <001>. We can notice that the FFT pattern for the substrate (Figure 5c) is similar and perfectly matches the one for the epilayer (Figure 5d), indicating the same structure. This observation confirms the epitaxial relationship between the epilayer and the substrate and clearly demonstrates the single-crystalline nature of the as-grown layers. As mentioned previously, the epilayer is oriented by the underlying substrate because the nucleation through the nanoholes is dominant. Similar observations have been reported by Yu et al. during the growth of GaN on nitrogen plasma-treated graphene.[36] The Ge epilayer's and the substrate's structures are diamond cubic as
<figure><img src="image_8.png" /><figcaption>**Figure 5.** TEM analysis of as-grown Ge layer on plasma-treated single-layer graphene (SLG) on Ge substrate using the APN approach. a) Cross-sectional HAADF-STEM image of Ge epilayer grown on the engineered SLG on Ge. The SLG interface is marked off by the red lines and the red arrow shows a small hole at the interface induced by the overgrowth on the SLG, b) cross-section HRTEM image of a defect-free sample (similar region displayed in Figure 4b) showing the graphene interface between the substrate and the epilayer. The yellow arrows mark off the regions where nucleation occurred directly on the substrate through the nanoholes created by the plasma treatment, c,d) FFT patterns of the epilayer and the substrate, respectively, showing the same diffraction spots demonstrating the epitaxial relationship between the substrate and the epilayer, and the monocrystalline nature of the as-grown epilayer, e,f) IFFT images showing the Ge(220) planes of the layer and the substrate highlighted by red arrows in (c) and (d), respectively, with a d-spacing of ≈0.2 nm, g–j) Geometrical phase analysis (GPA) deformation maps of the Ge layers grown on engineered SLG on Ge substrate converted from (b).</figcaption></figure>

shown in Figure 5f, displaying the inverted fast Fourier transform (IFFT) images that show the (220) planes of the epilayer and the substrate, respectively, with a d-spacing of ≈ 0.2 nm. Our results also demonstrated that some defects induced during the nucleation process due to multilayer graphene (see Figure 4a) can lead to the formation of a polycrystalline epilayer as shown in Figure S7 (Supporting Information) of the supplementary materials. From this figure, we can clearly see that the epilayer has random diffraction spots with different orientations compared to the substrate and there is evidence of stacking faults and microstructure twinning. This suggests that the initial graphene layer should be uniform and of high quality prior to any treatment.

To clarify the mechanisms of how the epilayers relieve the deformations and the epitaxial stress, the strain state of the as-grown Ge layers on engineered SLG was studied by HRTEM (Figure 5b) combined with the geometrical phase analysis (GPA)

method. The 2D strain and rotation maps of Ge/SLG/Ge are presented in Figure 5g–j. Note that the maps have been rotated by 90° with respect to the initial HRTEM image. The analysis of those maps provides clear evidence that there is no deformation and strains in the Ge epilayers with respect to the underlying substrate. This confirms that the growth is coherent without strain relaxation and comparable to homoepitaxy. From the in-plane map ($E_{xx}$), we can notice a very low negative out-of-plane deformation in the SLG interface that could be attributed to the nucleation on the graphene dangling bonds. However, in the regions we considered as nanoholes (Figure 5b), the in-plane deformation is zero ($E_{xx}$ = 0). These observations perfectly corroborate the mechanisms governing the APN approach that we presented in the previous sections. The quality of such layers is suitable for high-performance hybrid devices.

In comparison to RE,[11] which is governed by the polarity[12] of the substrate, and QVdWE governed by interactions with 2D
interface, the APN approach is a technique that can produce high-quality single-crystalline layers of semiconductors on any substrate including nonpolar substrate. Our findings provide new insights on plasma-induced defects in SLG and open a new route for the growth of hybrid functional devices and fabrication of semiconductor free-standing membranes. Few recent works reported on a similar technique using lateral overgrowth, where the nucleation is induced by defects or nanopatterns in the graphene.[18,24] For instance, a selective area epitaxy of 2D hBN on graphene induced by defects and the epitaxial MoS₂ on hBN assisted by substrate defects has been reported.[22,23] Our work demonstrates that defect-induced anchor points are promising for the epitaxial growth of any 3D semiconductors (polar and non-polar) on a graphene-terminated non-polar substrate, and alleviates the limitations of QVdWE and RE.

### 3. Conclusion

In this study, we report an original approach called anchor point nucleation for monolithic integration of 3D semiconductors on 2D materials. We demonstrated the epitaxial growth of defect-free single-crystalline Ge layers without any deformation on engineered SLG. Oxygen plasma treatments are used for surface engineering of the graphene, introducing defects such as dangling bonds and nanoholes in a uniform and controlled manner. These defects then act as preferential nucleation sites, enabling efficient and substrate-oriented nucleation for the growth of high-quality epilayers. HRTEM and STEM results provide clear evidence that the layers on SLG are perfectly single-crystalline and follow the underneath substrate orientation without any strain. Based on our experimental data, a growth model is proposed to elucidate the epitaxial growth mechanisms of Ge layers on engineered graphene. We demonstrated that the growth is governed by the nucleation through nanoholes followed by a lateral overgrowth of the nuclei anchored on the substrate and a coalescence of a continuous layer. These results lay the groundwork for advanced heterointegration of 3D/2D materials using engineered 2D substrates, for the development of highly mismatched structures and semiconductor free-standing membranes, as well as their disruptive applications.

### 4. Experimental Section

**Graphene Transfer:** Commercially available SLG was used in this work. The SLG grown by CVD on copper substrate was initially transferred on polymer support using polymethyl methacrylate (PMMA) layer and dissolution of the copper substrate. To transfer the SLG on the working substrate the received PMMA/SLG on a polymer support stack was slowly immersed in deionized water (DIW) until the PMMA/SLG layer was released from its support and was floating on top of the DIW. Before the transfer, the Ge (001) substrate was previously deoxidized in HBr for 1 min and rinsed with DIW. In fact, it was demonstrated that HBr could effectively remove suboxides on Ge surfaces.[21] This deoxidized substrate was introduced into the DIW and was used to carefully scoop out the PMMA/SLG from below. The Ge substrate with PMMA/SLG was then dried for 30 min in ambient clean room air, followed by annealing on the hot plate at 150 °C for 1 h. The sample was then stored for 24 h under the vacuum to avoid any detachment of the SLG and to remove all the residual moisture. Finally, to clean the SLG, the PMMA was dissolved in acetone and isopropyl alcohol (IPA) at 50 °C for 15 min each. This part was crucial to obtain a

PMMA-free SLG surface because any remaining polymer residue will result in unwanted nucleation leading to polycrystalline growth.[9]

**Plasma Treatment:** The SLG/Ge substrate was treated by an oxygen (O₂) plasma process to introduce defects in the SLG. The O₂ plasma treatment was performed in a barrel-type chamber, under a pressure of 320 mTorr and a radio frequency (RF) power of 10 W. The treatment duration varied between 6 and 60 s. It is worth mentioning that increasing the RF power significantly accelerates the etching process, leading to the complete destruction of the SLG. For example, at RF power equal to 20 W, the SLG was completely etched away after 12 s of the treatment. This aspect was carefully optimized during the experiments.

**Epitaxial Growth:** Epitaxial growth was carried out in VG Semicon VG90F CBE reactor, with liquid nitrogen cryopanel and with a thermocouple as a means of monitoring the temperature during the growth. Prior to epitaxial growth, all samples were annealed at 600 °C for 10 min to thermally deoxidize the revealed parts of Ge substrate and to help eliminate any potential residues of PMMA from the transfer. The Ge was grown at 520 °C, at a chamber pressure of ≈6E-6 Torr, using a solid source of Ge with a K-Cell temperature at 1250 °C.

**Characterizations:** Raman spectroscopy was performed using a Horiba Raman spectrometer with a CCD detector and a laser with an excitation wavelength of 474 nm to analyze defects induced in SLG by plasma treatment. A 100x objective was used for the measurements, which resulted in a laser spot of ≈1 µm in size. Before the data analysis, the background noise was subtracted. The D:G band ratio was calculated directly from the maximum intensity of the bands.

X-ray photoemission spectroscopy was realized with a Kratos Axis Ultra spectrometer using a monochromatic Al Kα source (hν = 1486.7 eV) and a charge-neutralized system. Survey scans and high-resolution scans were performed with an analysis area of 300 × 700 µm. The samples were loaded in the chamber, right after the plasma treatment, to avoid contaminations from ambient exposure. The XPS spectra were corrected to the main line of carbon 1s spectrum set to 284.8 eV. The data were analyzed using CasaXPS software. The Shirley background was subtracted from the XPS spectra, which were then fitted using a Voigt function.

The surface morphology of the nucleation and epitaxial layers was observed with a scanning electron microscope Zeiss LEO 1540 XB at 4.3 mm of working distance and 5 keV of acceleration voltage and with an atomic force microscope Veeco Dimension 3100 in tapping mode using SSS-NCHR silicon probe.

Prior to Scanning Transmission Electron Microscopy (STEM) and TEM imaging, the samples were prepared by focused ion beam (FIB) thinning and ion milling using a Zeiss NVision 40 Focused Ion Beam. The surface was protected by carbon coating to avoid the sputtering of the surface by ion beam. The TEM and STEM observations were made in Titan Themis microscope operated at 200 kV and equipped with a CEOS probe corrector and Ceta 16 M camera from FEI. The data were treated using Gatan digital micrograph software, to evaluate the crystalline quality and strains of the layers.

The crystalline quality of the epilayers was also investigated by the Rigaku smartlab HRXRD system with Cu Kα X-ray source in GIXRD configuration.

### Supporting Information

Supporting Information is available from the Wiley Online Library or from the author.

### Acknowledgements

T.M.D. and T.H. contributed equally to this work. The authors would like to thank Hubert Pelletier, Guillaume Bertrand, Philippe-Olivier Provost, and all the technical staff of 3IT for the technical support, Alexandre Chapotot for schematic drawing of the process, Javier Arias-Zapata and Bouraoui Ilahi for scientific discussions, Lilian Skokan for assistance during Raman spectroscopy measurements, and Sonia Blais for X-ray Photoelectron
spectroscopy measurements. The authors thank the Natural Sciences and Engineering Research Council of Canada (NSERC), Fonds de recherche du Québec (FRQNT), and Mitacs, for the financial support. A.B. is particularly grateful to the discovery grant for funding this research program. TEM studies were carried out within the C2N micro nanotechnologies platforms and partly supported by the RENATECH network and the General Council of Essonne. LN2 is a joint International Research Laboratory (IRL 3463) funded and co-operated in Canada by Université de Sherbrooke (UdeS) and in France by CNRS as well as ECL, INSA Lyon, and Université Grenoble Alpes (UGA). It is also supported by the Fonds de Recherche du Québec Nature et Technologies (FRQNT).

## Conflict of Interest

The authors declare no conflict of interest.

## Data Availability Statement

The data that support the findings of this study are available from the corresponding author upon reasonable request.

## Keywords

anchor point nucleation, graphene, heterointegration, substrate engineering, van der Waals epitaxy

Received: July 18, 2023

Revised: October 31, 2023

Published online: November 27, 2023

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