# Regular red-green-blue InGaN quantum wells with In content up to 40% grown on InGaN nanopyramids

Amélie Dussaigne, Colin Paillet, Névine Rochat, David Cooper, Adeline Grenier, Stéphane Vézian, Benjamin Damilano, Adrien Michon, Bérangère Hyot

► To cite this version:

Amélie Dussaigne, Colin Paillet, Névine Rochat, David Cooper, Adeline Grenier, et al.. Regular red-green-blue InGaN quantum wells with In content up to 40% grown on InGaN nanopyramids. Communications Materials, 2024, 5 (1), pp.280. 10.1038/s43246-024-00725-8 . hal-05219604

HAL Id: hal-05219604

https://hal.science/hal-05219604v1

Submitted on 22 Aug 2025

HAL is a multi-disciplinary open access
archive for the deposit and dissemination of sci-
entific research documents, whether they are pub-
lished or not. The documents may come from
teaching and research institutions in France or
abroad, or from public or private research centers.

L'archive ouverte pluridisciplinaire HAL, est
destinée au dépôt et à la diffusion de documents
scientifiques de niveau recherche, publiés ou non,
émanant des établissements d'enseignement et de
recherche français ou étrangers, des laboratoires
publics ou privés.

<figure><img src="image_1.png" /></figure>

Distributed under a Creative Commons Attribution 4.0 International License
<figure><img src="image_2.png" /></figure>

# Regular red-green-blue InGaN quantum wells with In content up to 40% grown on InGaN nanopyramids

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

Check for updates

Amélie Dussaigne<sup>1</sup>✉, Colin Paillet<sup>1,2</sup>, Névine Rochat<sup>1</sup>, David Cooper<sup>1</sup>, Adeline Grenier<sup>1</sup>,
Stéphane Vézian<sup>3</sup>, Benjamin Damilano<sup>3</sup>, Adrien Michon<sup>3</sup> & Bérangère Hyot<sup>1</sup>

Full color micro-displays with a pixel pitch of below 10 µm are needed for augmented and virtual reality applications. In the native emission approach, high efficiency Red-Green-Blue (RGB) pixels could be achieved using monolithically integrated InGaN based micro-LEDs. Here, we report the growth of high optical quality RGB InGaN/InGaN quantum wells grown on InGaN nanopyramids of diameter less than 1 µm by metal organic vapor phase epitaxy. We synthesized the nanopyramids by nanoselective area growth using an in situ patterned epitaxial graphene on SiC as an embedded mask. The RGB emission properties at different locations on the sample are dependent on the size of the InGaN nanopyramids. Advanced correlative analysis conducted on the same transmission electron microscopy lamella reveal a fully or at least nearly relaxed In<sub>0.13</sub>Ga<sub>0.87</sub>N core and very regular quantum wells emitting in the red range (620 nm) along the pyramid sidewalls with an In content up to 40%.

For virtual and augmented reality applications, micro-displays require high brightness Red-Green-Blue (RGB) sub-pixels of sizes below 10 × 10 µm<sup>21</sup>. The use of III-nitride materials promise high efficiency micro-light emitting diodes (micro-LEDs) compared to their organic counterparts. However, for such a pixel size, the pick and place process is no longer suitable for the combination of blue and green micro-LEDs from III-nitrides and red micro-LED from phosphide materials on the same platform. The mass transfer process is making good progress but the pixel pitch is still too large for such applications<sup>2</sup>. In addition, red emitting phosphide micro-LEDs suffer from an efficiency reduction with the reduced pixel size due to their long carrier diffusion length and to the presence of surface defects on micro-LED sidewalls<sup>3</sup>. A conversion approach can be used by depositing quantum dots or phosphors on blue micro-LEDs but their precise deposition is difficult and their resistance to luminous flow is low. Finally, having a single class of material family to get the three primary colors would enable native emissions and a monolithic integration of the RGB sub-pixels on the same substrate. To this aim, the In<sub>x</sub>Ga<sub>1-x</sub>N alloy is the best candidate as its bandgap can be tuned from blue to red emission by adjusting its In content. However, the external quantum efficiency (EQE) of III-nitride based LEDs drops for emission wavelength higher than 500 nm, to get only a few % in the red range<sup>4,5</sup>. This is mainly due to the material degradation for In content higher than 25% (inhomogeneous alloy, additional defects) when grown on GaN. Two main effects are at the origin of this degradation. First, the In<sub>x</sub>Ga<sub>1-x</sub>N bulk material is thermodynamically unstable for a high In

content, due to a large In covalent radius and a low In-N binding energy<sup>6</sup>. Second, the InGaN alloy is usually grown on GaN. It experiences in this case a strong compressive strain, which limits the In incorporation to an InN mole fraction of 25% by considering a coherent alloy<sup>7</sup>, for all growth conditions. Due to this strain state, In adatoms are pushed away from the GaN matrix during the growth of the InGaN layer to form a In reservoir on the surface. A part of this In reservoir can eventually be incorporated in the subsequent GaN layer. This effect is usually invoked as the compositional pulling effect<sup>8,9</sup>, or the In surface segregation effect<sup>10</sup>. In addition, on the c-plane of GaN, this compressive strain implies a strong piezoelectric polarization and thus a high internal electric field in the InGaN/(Ga,In)N quantum wells (QWs)<sup>11</sup>. This results in the quantum confined Stark effect, which causes a reduced transition energy and a spatial separation of the electron and hole wavefunctions for thick QWs. To get a red emission on c-plane GaN while not suffering from the internal electric field screening when increasing the carrier density, the InN mole fraction should reach 35% for a well width of 3 nm. The InN mole fraction would need to be even higher for a heterostructure grown on a semi-polar or non-polar plane. The important parameter we can control, to obtain such a high In content, is the reduction of the strain. Using a relaxed InGaN pseudo-substrate would be the best approach. By combining the full InGaN structure and the InGaNOS substrate from Soitec, an enhanced In incorporation rate with the increase of the in-plane lattice parameter of the substrate has been demonstrated<sup>12,13</sup>. InN mole fraction up to 39% was deduced in a particular area in between two V

<footnote n="1">¹Université Grenoble Alpes, CEA, Leti, 17 rue des Martyrs, F-38000 Grenoble, France. ²Now at Aixtron, Dornkaulstr. 2, 52134 Herzogenrath, Germany. ³Université Côte d'Azur, CNRS, CRHEA, Sophia Antipolis, rue Bernard Grégory, 06560 Valbonne, France. ✉ e-mail: amelie.dussaigne@cea.fr</footnote>
<figure><img src="image_4.png" /><figcaption>**Fig. 1 |** *Structure of the InGaN nanopyramids.* a Schematic representation of the structure, b HAADF/STEM image of a whole single nanopyramid emitting in the red range. The position of the graphene (Gr) layer is indicated by the orange line.</figcaption></figure>

shaped defects (V-pits) where the QWs were of high crystalline quality¹⁴. In the rest of the structure, the QWs were too inhomogeneous to get a reliable value of In content. A high compressive strain was probably present in these QWs to prevent at the same time a high In content and a high crystalline quality, owing to the likely high lattice mismatch between the in-plane lattice parameter of the buffer and the expected one for the bulk material of the red emitting thin QWs (i.e. well width of 2.5 nm with estimated In content ≥35%).

Different approaches are used in the literature to get efficient red emission. On GaN, two methods are distinguished. The first approach is the use of an AlGaN interlayer at the interface between the QW and the quantum barrier (QB) in order to reduce locally the strain¹⁵,¹⁶, to avoid In desorption during the QB growth¹⁷, and to modify the band structure to reduce the transition energy and increase the electron-hole wavefunction overlap⁵. The second one is to use the presence of V pits to create locally a strain reduction around the V pits and to improve the hole injection by the V pit sidewalls¹⁸,¹⁹. Combining one or the other approach with QWs emitting at different wavelengths in the active zone can conduct to a real improvement in the EQE of red LEDs with values of 5% at 630 nm for micro-LED size of 25 × 25 µm²²⁰, 10.5% at 625 nm for LED size of 1.39 × 1.40 mm²²¹, or to 16% at 620 nm for LED size of 1 cm²²¹. Good results on GaN on sapphire with EQE of 4.5% for a dominant wavelength of 616 nm for 6.7 × 6.7 µm² micro-LEDs have also been demonstrated²². As previously described, an InGaN pseudo-substrate can be employed to reduce the strain in the whole LED structure. For example, using porous GaN, an InGaN relaxation was achieved for 10×10 µm² InGaN/porous GaN mesas due to the compliance effect of the porous material. An EQE value of 0.2% at 8 A/cm² at 632 nm was obtained for 6 × 6 µm² red micro-LEDs²³. An EQE of 9% at 1 A/cm² for λ = 630 nm has also been demonstrated²⁴. In the case of InGaNOS, the EQE value was 0.14% at 8 A/cm² for λ = 625 nm in the case of 10 µm diameter micro-LED²⁵. This low value was explained at first order by a low light extraction efficiency (LEE) which was estimated below 4%. Finally, nanostructures such as nanowires²⁶ or pyramids²⁷,²⁸ are also a promising approach to reduce the strain in the whole InGaN based structure due to the relaxation from the free surfaces. An EQE of 9% at 630 nm has been achieved using axial nanowires²⁹. InGaN based pyramids²⁷ or truncated pyramids³⁰, both using a GaN seed layer, have been developed in the aim of offering a relaxed InGaN template. Red emission has been demonstrated in the case of InGaN based QWs grown on the InGaN/GaN pyramid sidewalls³¹ or in the case of InGaN based QWs grown on the c-plane facet of the truncated InGaN pyramid with an In content of 18%³² or 30%³³.

The growth of III-nitrides on 2D materials has attracted much attention to promote an easy transfer of the active structure on a host substrate³⁴,³⁵. 2D materials can also help to obtain relaxed III-nitride layers, such as a relaxed InGaN layer, due to the absence of dangling bonds on their

surface. However, the InGaN nucleation on graphene has been found to be very difficult³⁶. Using a different approach, graphene was used as a mask for selective area growth (SAG) with the advantage of connecting GaN to a high-electron-mobility and high-thermal-conductivity back electrode³⁷.

In this paper, we report the metal organic vapor phase epitaxy (MOVPE) SAG of InGaN hexagonal nanopyramids with an in situ patterned epitaxial graphene monolayer on SiC as an embedded mask. InGaN/InGaN multiple quantum wells (MQWs) are grown on the nanopyramid sidewalls. Depending on the size of the pyramid, an emission from blue to red, up to 645 nm at 300 K, is obtained on the same sample. For a nano-pyramid base diameter of 750 nm, the InGaN/InGaN QWs grown on the nanopyramid sidewalls exhibit a regular shape with an average thickness of 3.3 nm with a record InN mole fraction up to 40%. A homogeneous red emission centered at 620 nm at 10 K is measured by cathodoluminescence (CL) all along the sidewalls of the nanopyramids on the same transmission electron microscopy (TEM) lamella as the one used for TEM analysis. No piezoelectric polarization is observed in these heterostructures, confirming that the red emission arises from the high InN mole fraction.

## Results and discussion

Figure 1a shows a schematic of the structure. InGaN cores are grown on SiC through small apertures performed in situ in the epitaxial graphene layer. It is then followed by the growth of an InGaN based active zone on the nanopyramid sidewalls, which is composed of 5x InGaN/InGaN MQWs. Note that the use of a GaN seed layer was not necessary compared to previous works²⁷,³⁰. Figure 1b shows a low magnification image of a single InGaN nanopyramid obtained by TEM. It confirms the pyramidal shape and the position of the MQWs on the pyramid sidewalls. One can note that the MQWs are perfectly joined at the apex of the nanopyramid. The scanning electron microscopy (SEM) images in Fig. 2a, b show the pyramidal nuclei with an hexagonal base obtained after the MOVPE InGaN growth process on the graphene-on-SiC substrate previously prepared with an in situ annealing step. The nuclei appear preferentially organized along lines which follow the <1010> direction. This direction is typical of the orientation of the step edges of the SiC substrate (see Supplementary Fig. 1a). The graphene on the SiC step edges are more easily defective than on the terraces, and as such appear as favorable areas for the nucleation of the pyramids³⁸. Two opposite facets of each hexagon appear perpendicular to the direction of the step edges. An atomic force microscopy (AFM) profile passing through these two opposite facets and the apex of a pyramid indicates a facet inclination close to 61° (see Supplementary Fig. 1b). This inclination could correspond to the {1011} plane family whose theoretical angle with the growth direction <0001> is 61.96°. This is consistent with the alignment of the facets with the SiC <1010> direction (direction of the step edges (see Supplementary Fig. 1a)) observed on the SEM images. In
<figure><img src="image_5.png" /><figcaption>**Fig. 2** | Shape and orientation of the InGaN nanopyramids. a, b SEM images of nanopyramids at two different scales. c 2θ-ω XRD scan focused on the (0002) diffraction planes for the InGaN cores and the SiC substrate. d XRD φ-scan in the (1013) plane for InGaN and SiC.</figcaption></figure>

addition, it has been shown³⁹ that the use of H₂ in the carrier gas mixture results in etching of the {1011} planes and leads to the formation of pyramidal hexagonal-based nuclei. 2θ-ω X-Ray Diffraction (XRD) scans focused on the (0002) diffraction planes (Fig. 2c) reveal the presence of two peaks adjacent to each other, the first one located at 34.31° and the second one located at 34.38°. The presence of these two InGaN peaks could come from a slight demixing of InGaN in two different concentration phases, or from pyramids with different In contents. Considering a relaxed InGaN alloy, these positions would correspond to In concentrations of 5 and 7%, respectively. These concentrations are consistent with an expected In concentration of 6% in a planar InGaN layer grown in the same conditions. Figure 2d shows an in-plane XRD φ-scan acquired along the (1013) planes for InGaN and SiC. The presence of the six separated peaks for InGaN reveals that the pyramids have a common crystallographic alignment in the plane of the substrate. The superimposition of these peaks with the six SiC peaks demonstrates an epitaxial relationship between InGaN pyramids and SiC substrate.

Figure 3a presents the whole sample with a surface of 1 cm² under laser excitation (He-Cd laser emitting at 325 nm). An emission from blue to red is observed depending on the excited area. Photoluminescence (PL) measurements using a 405 nm laser diode show that the central emission wavelength ranges from 450 to 645 nm. Figure 3b exhibits the PL spectrum centered at 645 nm, taken in the red area on the sample. The linewidth of the spectrum is broad (105 nm) probably because of the variability of the size of the pyramids under the laser spot, whose diameter is estimated at 40 µm. By performing PL measurements as a function of the temperature (T) on the

same location, a plateau of PL integrated intensity vs 1/T was observed at low temperature until 60 K, before a drop of the PL integrated intensity (see Supplementary Fig. 2). The 15 K/290 K PL integrated intensity ratio is 3.3, which is good for red InGaN emission and is comparable, as an example, to what has been obtained previously on InGaNOS¹². Note that this does not constitute a direct evaluation of the internal quantum efficiency value but it gives an upper limit of it. CL measurements were then performed on single pyramids from different areas. Figure 4a-c presents SEM, and CL measurements of typical pyramids on three areas of the sample. The CL average spectra of the three pyramids present two contributions, the first one centered at 410 nm which is attributed to the InₓGa₁₋ₓN core, and the second one at higher emission wavelength which is attributed to the emission from the InₓGa₁₋ₓN/In₂Ga₁₋₂N MQWs. The contribution from the InₓGa₁₋ₓN core is consistent with an In content x of 13%. Then, the central emission wavelength of the second contribution shifts from 495 to 550 to 600 nm, when the base diameter of the pyramid increases from 400 to 700 to 1000 nm, respectively. The variation of the pyramid size across the sample is also accompanied by a variation of the pyramid density. Lower density leads to the formation of bigger pyramids. The available total InGaN surface is lower in areas where the pyramids are bigger, by considering the pyramid sidewall total surface and the nanostructure density. It leads to a higher InGaN growth rate during the QW growth in the red emitting area as the In and Ga adatoms are distributed between the available semi-polar surfaces of the InGaN cores. It results in thicker well width and higher In content⁴⁰ in this region of the sample, explaining the longer emission wavelength. The emission wavelength depends thus on the size of the pyramids and their
<figure><img src="image_6.png" /><figcaption>**Fig. 3** | Emission at the macro scale of InGaN nanopyramids from different areas on the sample. a Pictures of the sample under a 325 nm laser excitation at different positions on the sample, b PL spectrum at RT in the red area with a 405 nm laser excitation (corresponding to the last picture on the right of (a)).</figcaption></figure>

<figure><img src="image_7.png" /></figure>

<figure><img src="image_8.png" /><figcaption>**Fig. 4** | Emission at the nanoscale of single InGaN nanopyramids of different size. a–c SEM image of a single nanopyramid of base diameter of 500 nm, 750 nm and 1000 nm, respectively. d–f CL average spectrum on this single pyramid of base diameter of 500 nm, 750 nm and 1000 nm, respectively. g–i CL integrated intensity</figcaption></figure>

mapping in the QW wavelength range (450–750 nm) of this single nanopyramid of base diameter of 500 nm, 750 nm and 1000 nm, respectively. The dotted black line refers to the position of the InGaN core contribution and the red line refers to the central emission wavelength of the QW emission.
<figure><img src="image_9.png" /><figcaption>**Fig. 5 |** Optical and structural properties in the QW area of single InGaN nanopyramids of different size. a–c SEM image of a TEM lamella containing a single nanopyramid of base diameter of 500 nm, 750 nm and 1000 nm, respectively. d–f 10 K CL average spectrum calculated in the white delimited area on the SEM image of this single nanopyramid of base diameter of 500 nm, 750 nm and 1000 nm,</figcaption></figure>

respectively. g–i CL wavelength mapping of the maximum intensity QW peak position of this single nanopyramid of base diameter of 500 nm, 750 nm and 1000 nm, respectively. j–l HRSTEM image in the QW area performed on the same TEM lamella of these nanopyramids of base diameter of 500 nm, 750 nm and 1000 nm, respectively.

density in each area of the sample. The CL MQW integrated intensity images reveal for these three examples that the CL intensity is homogeneous from sidewall to sidewall of each pyramid; and that the CL intensity is increased at the pyramid ridges and at the apex, as expected due to the enhance extraction efficiency on these locations. Note that for the 400 nm pyramid, the pyramid shape on the CL intensity mapping is disturbed due to a specimen drift during the low temperature measurement that is more pronounced during smaller size object mapping (see Supplementary Fig. 3).

To understand more deeply these structures, three TEM lamella were prepared by focus ion beam (FIB). These three TEM lamellas are composed of one single pyramid with base diameters of 500, 750 and 1000 nm, respectively. The low temperature (10 K) CL mappings performed on the cross section of these TEM lamellas are exhibited on Fig. 5. The central emission wavelength of the average CL spectra switches from 550 nm to 597 to 618 nm, for pyramid base diameters of 500, 750, and 1000 nm, respectively. Their full width at half maximum (FWHM) varies between 54, 58, and 64 nm, respectively. The side peak on the left side of each spectrum is related to the emission of the InGaN core. The QW emission is well located along the pyramid sidewalls. Except for the left edge of the 1 µm pyramid, the emission wavelength is quite homogeneous along the sidewalls of the three different pyramids. High resolution scanning TEM (HRSTEM) images are also provided for each type of pyramid, and were recorded on the exact same pyramid. The average well width is 2.2 nm, 3.2 nm, and 4.5 nm for 500 nm, 750 nm and 1000 nm pyramid size, respectively. No thickness variation along each QW is observed in these images. On the other hand, the QW (QB) width decreases from the first QW (QB) on the core side of the pyramid to the last QW (QB) near the surface, for each type of pyramid. This is due to the increased diameter of the pyramid after the growth of each layer of the active area, increasing thus the available growth surface of the pyramid before the growth of the next layer. Indeed, the material (i.e. In and Ga adatoms) is distributed over the total available surface (here the surface of the semi-polar planes of the pyramid sidewalls). If this total available surface

increases, then the growth rate of the next layer is decreased, leading to a reduced well or barrier width. The different thicknesses in the case of the 750 nm pyramid are given in Supplementary Fig. 4 as an example. These different values of well and barrier widths can also explain the broad average CL spectrum while the peak emission wavelength seems homogeneous along the pyramid sidewalls. To reduce the linewidth, the increased pyramid diameter during the growth of the active zone should be taken into account.

A closer look is now taken on the 750 nm nanopyramid emitting in the red range. Figure 6a shows that, at low magnification, the nanostructure has a perfect pyramidal shape. The QWs seem homogeneously deposited along the sidewalls, and they perfectly join at the apex. First, the InGaN core presents some extended defects just above the hole in the graphene. One vertical defect goes through the whole height of the pyramid to finish at its apex. Two other extended defects cross the pyramid horizontally from the middle to the sidewall at the bottom or at half of the height of the pyramid. The rest of the pyramid seems free of extended defects from this orientation. The images from Fig. 6b, c show the interface between the InGaN core and the graphene, and between the InGaN core and the SiC substrate, respectively. The interfaces seem abrupt in both cases. The graphene layer is well visible and shows the van der Waals (vdW) gap provided by this 2D material. The hole in the graphene has a diameter of 100 nm (see Supplementary Note 1 for hole diameter of the other pyramids). Energy dispersive X-ray spectroscopy (EDX) performed at the base of the pyramid, above the aperture and above the graphene, reveals an In content of 13% in both positions. Another EDX profile was taken from the base of the pyramid above the aperture and then vertically through the whole pyramid until its apex (Fig. 6d, e). In this case also, the EDX profile shows a plateau of In content at the value of 13%. The In content is constant from the base to the apex with an increase only in the QW area. Finally, the whole InGaN core presents an In content of 13%, which could mean that the InGaN pyramid core is at least nearly relaxed. This is quite surprising since usually a strain gradient is observed from the base to the apex and along the sidewalls of a
<figure><img src="image_10.png" /><figcaption>**Fig. 6** | Structural and chemical properties of the InGaN core from the red emitting InGaN nanopyramid. a Bright field STEM image of the whole red emitting nanopyramid. HAADF HRSTEM images at the graphene/InGaN interface and in inset the corresponding EDX mapping in the InGaN core (b) on the left side of the</figcaption></figure>

nanopyramid base, c near the opening on the SiC substrate. d EDX mapping of the whole nanopyramid, e EDX profile taken from the base to the apex of the nano-pyramid from the blue rectangle of (d), the blue arrow in both figures gives the profile direction. f Strain mapping of the same nanopyramid along the c axis.

<figure><img src="image_11.png" /><figcaption>**Fig. 7** | Structural and optical properties of the red emitting active zone. a HAADF HRSTEM image of the red emitting nanopyramid in the QW area. The white arrow indicates the growth direction. CL measurements on the same TEM</figcaption></figure>

lamella with (b) SEM image of the whole nanopyramid, c CL peak wavelength mapping, d CL intensity mapping, and e CL FWHM mapping.

pyramid, leading to an emission wavelength gradient from the QWs grown on the sidewalls with an emission redshift from the base to the apex of the pyramid²⁸. Other groups²⁷,³⁰,³³,⁴¹ observed also an In content gradient from the base to the apex of the core of their InGaN pyramids, which could be truncated or not. Note that in these cases the base of the pyramid consists of a GaN seed layer which has been found essential to get perfect pyramid shape²⁷, and that the mask used for the SAG is SiₓNᵧ. The pyramids are thus not truly full InGaN pyramids compared to our current case. Here, the

InGaN core is probably nearly relaxed already from the base of the nano-pyramid due to the high density of defects present above the hole in graphene, formed during the nucleation step on SiC. The role of the graphene is at least to favor the adatom diffusion length in between two pyramids, as no dangling bonds are present on its surface. It is especially of high importance for the In adatoms to promote their diffusion towards the pyramid sidewalls before their desorption, in order to incorporate a high amount of In in the pyramid core and in the QWs. Another role of the graphene could be to
<figure><img src="image_12.png" /><figcaption>**Fig. 8** | Deformation and chemical mappings of the red emitting active zone. a HRSTEM image in the QW area on the same TEM lamella, the white arrow indicates the growth direction. b Corresponding deformation mapping along the {1011} axis obtained by GPA, with in inset a profile taken at the top of the deformation mapping. c EDX mapping on the same TEM lamella of the red emitting nanopyramid for Ga (green) and In (pink), d EDX profile from the blue rectangle of (c), the blue arrow in both figures indicates the profile direction. The red dashed line in (d) indicates the position of the In concentration for a value of 40%.</figcaption></figure>

promote the relaxation of the InGaN core from the base of the pyramid. Indeed, as the hole diameter is small compared to the base diameter of the pyramid, one can suggest that the vdW gap between graphene and InGaN (i.e. in regions where parts of the InGaN core are above the graphene) plays a role on the strain relaxation at the early stage of the pyramid growth. Further investigations should be performed to answer this question, such as for example growing same InGaN nanopyramids through a SiₓNᵧ mask on SiC with same hole diameter and density.

To provide a high precision deformation map across a large area, precession electron diffraction (PED) was used⁴². Diffraction patterns were acquired to provide a map with a field of view of 800 × 800 nm² with a step size of 8 nm and a spatial resolution given by the probe size of around 2 nm. The displacements of the diffracted beams are then compared to a reference to provide a deformation map. Figure 6f shows a deformation map of this nanopyramid along the *c* axis. The reference is taken in the middle of the InGaN core at half of the pyramid height, so that we are only looking at the variation in the core and the measured deformation values are independent of the strain generated by the InGaN growth on SiC inside the aperture. From the whole core, only a very small deformation gradient is observed with a variation of displacement of only +0.5% ±0.05 %. A strain variation is indeed present between the area of the aperture and just above (area 1 in blue) and the rest of the pyramid (area 2 in white). This low strain variation inside the core accompanied with a homogeneous In content of 13% well support the hypothesis of a fully or at least nearly relaxed InGaN core, at least in area 2, as the strain at the apex of a pyramid should be relaxed, and as the In incorporation is strongly dependent of the strain state.

Figure 7a shows the high angle annular dark field (HAADF)-HRSTEM image in the QW area. In this high magnification image, the well width decreases from the core to the sidewalls, i.e. from the first QW to the fifth QW, as explained above, but the thickness of each QW is regular along the sidewalls and their chemical contrast is homogeneous. The same trend is observed for the QBs. Here the well width is reduced

from 3.6 to 2.6 nm from the core to the surface of the pyramid sidewall (well width of 3.6 nm, 3.6 nm, 3.1 nm, 3.1 nm, and 2.6 nm from the first to the fifth QW). The barrier width decreased as well from the core to the pyramid surface from 8.2 nm to 6.7 nm (barrier width of 8.2 nm, 9 nm, 6.7 nm, 7.2 nm). For clarity, these thicknesses were reported on the TEM image in Supplementary Fig. 4. Figure 7c, d, e presents the CL mappings (peak wavelength, intensity and FWHM mappings, respectively) recorded on the same TEM lamella. The emission is also homogeneous along the sidewalls, from the base to the apex, and is ranging from 590 to 625 nm with an average emission wavelength centered at 618 nm. Only a part of the sidewalls on the left side of the mappings is not emitting, probably due to the presence of an extended defect similar to that observed on Fig. 6a. The CL intensity mapping well follows the CL peak wavelength one with a homogeneous CL intensity in the red emitting areas on the sidewalls. The FWHM mapping reveals values from 50 to 90 nm with an average linewidth of 77 nm. This broad emission could be explained by the well width decrease from the core to the sidewalls, as explained in the previous paragraph. Examples of local spectra from the top left or the bottom right of this nanopyramid can be found in Supplementary Fig. 5. As the QWs are of good crystalline quality, it was possible to measure the deformation at high resolution in the QWs using the geometrical phase analysis (GPA) technique⁴³. Here a Fourier transform is applied to the HRSTEM image in Fig. 8a and information about the lattice planes in the {1011} growth direction can be selected to calculate the deformation map shown in Fig. 8b. By taking the reference from the InGaN core, a deformation value of 5.5% ±1 % is obtained in the QWs along the direction perpendicular to the surface of the sidewalls. Such a high deformation in the QWs should be caused by a high In content, especially starting from an In₀.₁₃Ga₀.₈₇N buffer which is taken at the origin of the deformation. EDX measurements were then performed on these QWs and quantified using Cliff-Lorimer method¹⁴. Figure 8c shows the EDX mapping in the QW area and Fig. 8d presents an EDX
**Fig. 9 | Electron holography measurements on the red emitting active zone.** a Potential map of the red emitting nanopyramid on the same TEM lamella as all the other nanoscale measurements, b expanded view of the potential map of (a) corresponding to the white area of (a). For specimens grown along the (0001) direction a negative to positive potential gradient is expected from the base to the top of each QW. This is not observed here. c Corresponding amplitude image showing that the region of interest is free from diffraction contrast. d Potential profile averaged across the region shown in (b).

<figure><img src="image_13.png" /></figure>

profile obtained from an averaged EDX profile. The InN atomic fraction in the QBs is of 8% and is higher than 40% in the QWs.

To complete our study, off-axis electron holography was also performed on the same TEM lamella in order to measure the piezoelectric potential in the QW zone. Figure 9a displays a reconstructed phase image and Fig. 9b shows an expanded view of the region indicated in (a). The phase of the electrons is sensitive to the potential and as such we can measure the polarization in the nano-pyramid. The specimen is carefully tilted to reduce the diffraction contrast while keeping the QWs on the left hand side of the pyramid aligned with the electron beam. The reconstructed amplitude image presented in Fig. 9c shows that the diffraction contrast has been reduced and the phase image can indeed be interpreted in terms of the potential in the specimen. The corresponding calculated electric potential profile is shown in Fig. 9d. Usually, for QWs grown along the c-plane direction, two contributions are visible in the phase image: the mean inner potential which is a property of the material and thus directly related to the In content in the case of InGaN, and the polarization potential¹⁵. In specimens grown in the (0001) direction, a strong piezoelectric field would be expected which is observed as a negative to positive potential gradient across the QWs. However, here only the mean inner potential contribution is observed. Thus the strong positive piezoelectric effect has been reduced in these QWs grown along the {1011} direction as predicted elsewhere⁴⁶. In general, due to the high sensitivity of the technique, it is straightforward to measure typical values of the piezoelectric potential in InGaN QWs. As such we can conclude that the internal electric field is thus very low in these heterostructures. This is confirmed with PL measurements as a function of the excitation power: a wavelength shift of only 14 nm is measured with PL excitation power densities from 0.88 to 240 W.cm⁻² (see Supplementary Fig. 6).

Band structure simulations were conducted using the Silense software from STR⁴⁷. The exact structure described in this article (different QW and QB thicknesses) was implemented using 13% of InN mole fraction in the buffer layer, 8% in the quantum barriers, 45% for the first QW, 43% for the

fourth QW, and 42% for the others, as measured by EDX (maximum values in the QWs from Fig. 8d). The buffer is supposed to be fully relaxed. The orientation of the structure is defined with an angle of 60° from the c-plane. The band structure obtained together with the wave functions of each QW are displayed in Fig. 10a. An average internal electric field of 0.8 MV.cm⁻¹ is found. It is not zero as supposed from the e-holography results but is very low anyway for an average In content of 42% in the QWs. Indeed, it should be between 5 and 6.6 MV.cm⁻¹ on c-plane for such a high InN mole fraction in the QWs¹¹,⁴⁸. The emission spectrum is obtained from the simulation of a 380 nm laser excitation with a power density of 1 W.cm⁻² at 300 K (Fig. 10b). The shape of the spectrum could be at first surprising but it actually accounts for the emission of each of the five QWs with different well width and/or InN mole fractions, leading to different contributions in the PL spectrum. The peak emission wavelength is 632 nm. The linewidth is 85 nm. It is well in agreement with the PL and CL results. The linewidth is similar to what is obtained by experiment without the need to take into account Inhomogeneities while it is usually the case for the InGaN alloy⁴⁹. Indeed, this broad linewidth is the resultant of the five QWs with different thicknesses and In content. Note that the linewidth could be reduced by a factor 4 with QWs of same thickness according to the simulation using average In content and thickness values for QWs and QBs.

This high InN mole fraction of 42% in the QWs can be obtained through the use of a relaxed or nearly relaxed In₀.₁₃Ga₀.₈₇N core, compared to the InN mole fraction limit of 25% when grown on GaN, and thus when the InGaN alloy is in a strong compressive strain. Compared to our previous results on the InGaNOS substrate from Soitec, the InN mole fraction is indeed higher with values higher than 40%. Furthermore, the In₃Ga₁₋₃N/In₂Ga₁₋₂N QWs are more homogeneous in terms of thickness and chemical contrast. This would mean that a strong reduction of the strain in the active zone helps not only to increase the In incorporation rate but also provides a more homogeneous coherent In₃Ga₁₋₃N alloy. In the current case, the in-plane lattice parameter *a* of the relaxed or nearly relaxed In₀.₁₃Ga₀.₈₇N core should be not far from the target of 3.238 Å (i.e. relaxed In₀.₁₄Ga₀.₈₆N layer)
<figure><img src="image_14.png" /><figcaption>**Fig. 10** | Band structure simulations of the red emitting active zone. a Band diagram and (e-h) wave functions simulated with Silense software of the red emitting MQW structure with various well and barrier thicknesses. b Corresponding simulated PL spectrum with a 380 nm excitation at 1 W.cm⁻².</figcaption></figure>

which would place the red emitting QWs in the same strain state as blue QWs grown on GaN⁵⁰. Note in addition that these structures are free of V shaped defects and present a very low dislocation density in the active zone. The pyramid configuration should provide in addition a high light extraction efficiency.

## Conclusion

In conclusion, red emitting InGaN nanopyramids grown directly on SiC through apertures in a thin layer of graphene provide regular InGaN/InGaN quantum wells with an extremely high In content of at least 40%. A low internal electric field is revealed both by electron holography and simulations which confirm such a high InN mole fraction and promise the ability of this structure to keep red emission whatever the injected carrier density. The graphene layer used as a mask promotes the diffusion of adatoms to get high In content in the pyramid core and in the QWs grown on the sidewalls. Blue to red emission was also demonstrated in one growth run at different locations on the sample. The emission wavelength depends on the nano-pyramid diameter and density. This is a promising approach for monolithically integrated RGB micro-LEDs for micro-display applications. To go further, it will necessitate to define an organized pattern in the graphene layer and to complete the LED structure by a p-doped InGaN layer grown on the InGaN nanopyramid sidewalls. Then, the full integration for micro-display based devices will probably imply the transfer of the InGaN nanopyramids on a host substrate. This approach could be implemented also on

Si substrate with a transfer of a thin graphene layer. Finally, from a more general point of view, these results demonstrate that with an appropriate reduction of the strain in the active zone, a very high In content can be incorporated in InGaN QWs without degrading the InGaN alloy.

## Methods

### Epitaxy of graphene

The first step in our process is the epitaxial growth of graphene on the Si-face of 6H-SiC substrates through propane-hydrogen chemical vapor deposition (CVD)⁵¹. The SiC substrates have a 0.1° miscut angle. The growth performed at 1550 °C at a pressure of 800 mbar, with Ar, H₂ and C₅H₈ fluxes of 10 slm, 2 slm and 10 sccm respectively, for 5 min results in the formation of an atomically flat graphene layer over the entire substrate stepped surface⁵². After the complete coverage of the surface by the monolayer (ML) graphene, no additional growth occurs indicating a self-limiting process⁵¹. The ML graphene is actually lying on a carbon layer partially bonded to the SiC substrate, called the buffer layer. Raman spectrum of the epitaxial graphene displays the well-established 2D ($\omega_{2D} = 2686$ cm⁻¹) and G ($\omega_G = 1588$ cm⁻¹) peaks after the subtraction of the SiC background. Both G and 2D peaks show narrow full width half maxima of 9 and 24 cm⁻¹, respectively. No D peak (expected around 1350 cm⁻¹) related to the defects was observed. These features demonstrate a good structural quality of the graphene layer.

### Epitaxy of the InGaN nanopyramids

The sample was then transferred for MOVPE of III-nitrides in a close-coupled showerhead 6 × 2 inch Aixtron reactor. An in situ preparation step of the graphene layer at 1100 °C under NH₃ and H₂ ambient allows to locally deteriorate the graphene layer in order to create randomly distributed holes with a diameter of around 100 nm. Then, InGaN pyramids were grown by nucleating first on the SiC substrate through the graphene layer. The InₓGa₁₋ₓN cores were grown at a growth temperature of 800 °C under a N₂/H₂ gaz mixture. A small amount of H₂ (5%) was indeed necessary to obtain faceted islands with a hexagonal base. The intended In content x was 6%, by comparison to a planar InGaN layer grown in the same growth conditions. Then, 5x InₓGa₁₋ₓN/In₂Ga₁₋₂N MQWs were grown with the same growth conditions as those used to get red emission on the InGaNOS substrate⁵³. Only the growth time was adjusted to account for the limited total available growth surface on the InGaN cores. The structure is finished by a 30 nm thick In₂Ga₁₋₂N layer. A scheme of the structure and a STEM image of a single nanopyramid are provided on Fig. 1a, b, respectively.

## Material characterization

The sample was then characterized by SEM, XRD, PL at room temperature using an He-Cd laser or a 405 nm laser diode with a laser spot of 100 µm or 40 µm diameter, respectively, and low temperature (10 K) CL with 3 keV electron acceleration voltage on top view single pyramids. Three TEM lamella, each including an InGaN nanopyramid with a different diameter were prepared by FIB. Finally, correlative characterizations were conducting on the same TEM lamella, such as HRSTEM, low temperature (10 K) CL with 10 kV acceleration voltage, EDX, PED, and electron holography. Special care was taken during the FIB preparation to reduce amorphization of the pyramid cross section surface. The HRSTEM images of the interface between the SiC, graphene and InGaN pyramids were acquired at 80 kV on the FEI Titan Themis to avoid damaging the specimen. Further HAADF STEM images of the InGaN QWs were then acquired at 200 kV. The EDX was performed on the FEI Titan Themis operated at 200 kV. For quantification the 3 keV In edge and 9 keV Ga edges were used in a Cliff Lorimer method. The composition maps were then compared to a well known calibration specimen to assess the accuracy of the quantification. Precession electron diffraction was performed on a probe corrected FEI Themis operated at 200 kV. The precession is used to remove the dynamical effects in the diffraction patterns which leads to increased precision. The spatial resolution of 8 nm was given by the step size of the acquisition over the large field of view rather than the convergence angle of the measurement. The acquired diffraction patterns were analyzed using dedicated home made
software at CEA. For the electron holography, an image corrected FEI Titan
Ultimate was used operated at 200 kV. To achieve high sensitivity, stacks of
32 electron holograms were acquired of the specimen each recorded for 8 s.
The fringe spacing was 1.5 nm to give a nominal spatial resolution of 3 nm in
the reconstructed phase images. Data sets were acquired at many specimen
tilts to assess the accuracy of the measurements. Care was taken to make sure
that the QWs were aligned with the direction of the electron beam. These
experimental settings have been shown to easily image the piezoelectric
potentials in c-plane InGaN QWs.

Data availability

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

Received: 4 July 2024; Accepted: 18 December 2024;
Published online: 27 December 2024

References

1. F. Olivier et al. Influence of size-reduction on the performances of GaN-based micro-LEDs for display application. *J. Lumin.* **191**, 112–116 (2017).

2. Liu, Y.-T., Lin, C.-L., Lin, T.-Y., & Li, Y.-L. Performance of MicroLED Chip and display for emerging applications. ICNS14, OD15-1 (2023)

3. Li, P. et al. Size-independent peak external quantum efficiency (>2%) of InGaN red micro-light-emitting diodes with an emission wavelength over 600nm. *Appl. Phys. Lett.* **119**, 081102 (2021).

4. Auf der Maur, M., Pecchia, A., Penazzi, G., Rodrigues, W. & Di Carlo, A. Efficiency drop in green InGaN/GaN light emitting diodes: The role of random alloy fluctuations. *Phys. Rev. Lett.* **116**, 027401 (2016).

5. Hwang, J.-I., Hashimoto, R., Saito, S. & Nunoue, S. Development of InGaN-based red LED grown on (0001) polar surface. *Appl. Phys. Express* **7**, 071003 (2014).

6. Ho, I. & Stringfellow, G. B. Solid phase immiscibility in GaInN. *Appl. Phys. Lett.* **69**, 2701 (1996).

7. Lymperakis, L. et al. Elastically frustrated rehybridization: Origin of chemical order and compositional limits in InGaN quantum wells. *Phys. Rev. Mat.* **2**, 011601 (2018).

8. Pereira, S. et al. Compositional pulling effects in InₓGa₁⁻ₓN/GaN layers: A combined depth-resolved cathodoluminescence and Rutherford backscattering/channeling study. *Phys. Rev. B* **64**, 205311 (2001).

9. Inatomi, Y. et al. Theoretical study of the composition pulling effect in InGaN metalorganic vapor-phase epitaxy growth. *Jpn. J. Appl. Phys.* **56**, 078003 (2017).

10. Dussaigne, A., Damilano, B., Grandjean, N. & Massies, J. In surface segregation in InGaN/GaN quantum wells. *J. Cryst. Growth* **251**, 471 (2003).

11. Bernardini, F., Fiorentini, V. & Vanderbilt, D. Spontaneous polarization and piezoelectric constants of III-V nitrides. *Phys. Rev. B* **56**, R10024 (1997).

12. Even, A. et al. Enhanced In incorporation in full InGaN heterostructure grown on relaxed InGaN pseudo-substrate. *Appl. Phys. Lett.* **110**, 262103 (2017).

13. Cooper, D., Boureau, V., Even, A., Barbier, F. & Dussaigne, A. Determination of the internal piezoelectric potentials and indium concentration in InGaN based quantum wells grown on relaxed InGaN pseudo-substrates by off-axis electron holography. *Nanotechnology* **31**, 475705 (2020).

14. Samuel, B. et al. Origins of nanoscale emission inhomogeneities of high content red emitting InGaN/InGaN quantum wells. *J. Appl. Phys.* **129**, 173105 (2021).

15. Lekhal, K. et al. Strain-compensated (Ga,In)N/(Al,Ga)N/GaN multiple quantum wells for improved yellow/amber light emission. *Appl. Phys. Lett.* **106**, 142101 (2015).

16. Ewing, J. J. et al. High external quantum efficiency (6.5%) InGaN V-defect LEDs at 600 nm on patterned sapphire substrates. *Speck Opt. Express* **31**, 41351 (2023).

17. Iida, D., Zhuang, Z., Kirilenko, P., Velazquez-Rizo, M. & Ohkawa, K. Demonstration of low forward voltage InGaN-based red LEDs. *Appl. Phys. Express* **13**, 031001 (2020).

18. Zhang, S. et al. Efficient emission of InGaN-based light-emitting diodes: toward orange and red. *Photonics Res* **8**, 1671 (2020).

19. Marcinkevičius, S. et al. Experimental evidence of hole injection through V-defects in long wavelength GaN-based LEDs. *Speck Appl. Phys. Lett.* **123**, 201102 (2023).

20. Huang, Y.-M. et al. High-efficiency InGaN red micro-LEDs for visible light communication. *Photonics Res* **10**, 1978 (2022).

21. Lee, D. et al. High-efficiency InGaN red light-emitting diodes with external quantum efficiency of 10.5% using extended quantum well structure with AlGaN interlayers. *Appl. Phys. Lett.* **124**, 121109 (2024).

22. Armitage, R. et al. True-red InGaN light-emitting diodes for display applications. *Phys. Status Solidi RRL* 2400012. https://doi.org/10.1002/pssr.202400012 (2024).

23. Pasayat, S. S. et al. Demonstration of ultra-small (<10 µm) 632 nm red InGaN micro-LEDs with useful on-wafer external quantum efficiency (>0.2%) for mini-displays. *Appl. Phys. Express* **14**, 011004 (2021).

24. Zhu, T. Native InGaN RGB: Empowering display applications with full color microLEDs, Proceedings Volume PC12886, Gallium Nitride Materials and Devices XIX; PC1288612. https://doi.org/10.1117/12.3003668 (2024).

25. Dussaigne, A. Full InGaN red (625 nm) micro-LED (10 µm) demonstration on a relaxed pseudosubstrate. *Appl. Phys. Express* **14**, 092011 (2021).

26. Zhang, X. et al. Chemical composition fluctuations and strain relaxation in InGaN nanowires: The role of the metal/nitrogen flux ratio. *Mat. Sci. Semicond. Process.* **55**, 79 (2016).

27. Bi, Z. et al. High In-content InGaN nano-pyramids: Tuning crystal homogeneity by optimized nucleation of GaN seeds. *J. Appl. Phys.* **123**, 025102 (2018).

28. Bi, Z. et al. InGaN platelets: Synthesis and applications toward green and red light-emitting diodes. *Nano Lett.* **19**, 2832 (2019).

29. Pandey, A. et al. A red-emitting micrometer scale LED with external quantum efficiency >8%. *Appl. Phys. Lett.* **122**, 151103 (2023).

30. Cai, W. et al. High In content nitride sub-micrometer platelet arrays for long wavelength optical applications. *Appl. Phys. Lett.* **121**, 211105 (2022).

31. Ko, Y.-H. et al. Red Emission of InGaN/GaN Double Heterostructures on GaN Nanopyramid Structures. *ACS Photonics* **2**, 515 (2015).

32. Bi, Z., Gustafsson, A. & Samuelson, L. Bottom-up approaches to microLEDs emitting red, green and blue light based on GaN nanowires and relaxed InGaN platelets. *Phys. B* **32**, 018103 (2023).

33. Cai, W. et al. Red emission from InGaN active layer grown on nanoscale InGaN pseudosubstrates. *J. J. Appl. Phys.* **62**, 020902 (2023).

34. Shin, J. et al. Vertical full-colour micro-LEDs via 2D materials-based layer transfer. *Nature* **614**, 81 (2023).

35. Bayram, C. et al. Vertical thinking in blue light emitting diodes: GaN-on-graphene technology. Dimitrakopoulos Proc. SPIE 9364, Oxide-based Materials and Devices VI, 93641C (2015).

36. Paillet, C. et al. InGaN islands and thin films grown on epitaxial graphene. *Nanotechnology* **31**, 405602 (2020).

37. Puybaret, R. et al. Nanoselective area growth of GaN by metalorganic vapor phase epitaxy on 4H-SiC using epitaxial graphene as a mask. *Appl. Phys. Lett.* **108**, 103105 (2016).

38. Robinson, J. et al. Nucleation of Epitaxial Graphene on SiC(0001). *ACS Nano* **4**, 153 (2010).

39. Lai, S. et al. Polarity and Its Influence on Growth Mechanism during MOVPE Growth of GaN Sub-micrometer Rods. *Cryst. Growth Des.* **11**, 1573 (2011).
40. Bosi, M. & Fornari, R. A study of Indium incorporation efficiency in InGaN grown by MOVPE. *J. Crys. Growth* **265**, 434 (2004).

41. Gustafsson, A., Bl, Z. & Samuelson, L. From InGaN pyramids to micro-LEDs characterized by cathodoluminescence. *Nano Express* **2**, 014006 (2021).

42. Cooper, D., Denneulin, T., Bernier, N., Béché, A. & Rouvière, J.-L. Strain mapping of semiconductor specimens with nm-scale resolution in a transmission electron microscope. *Micron* **80**, 145 (2016).

43. Hýtch, M. J., Snoeck, E. & Kilaas, R. Quantitative measurement of displacement and strain fields from HREM micrographs. *Ultramicroscopy* **74**, 131 (1998).

44. Watanabe, M. & Williams, D. B. The quantitative analysis of thin specimens: a review of progress from the Cliff-Lorimer to the new ζ-factor methods. *J. Microsc.* **221**, 89 (2006).

45. Boureau, V. & Cooper, D. Highly spatially resolved mapping of the piezoelectric potentials in InGaN quantum well structures by off-axis electron holography. *J. Appl. Phys.* **128**, 155704 (2020).

46. Romanov, A. E., Baker, T. J., Nakamura, S. & Speck, J. S. Strain-induced polarization in wurtzite III-nitride semipolar layers. *J. Appl. Phys.* **100**, 023522 (2006).

47. https://str-soft.com-devices/silense/.

48. Lefebvre, P. et al. High internal electric field in a graded-width InGaN/GaN quantum well: Accurate determination by time-resolved photoluminescence spectroscopy. *J. Massies Appl. Phys. Lett.* **78**, 1252–1254 (2001).

49. Piccardo, M. et al. Localization landscape theory of disorder in semiconductors. II. Urbach tails of disordered quantum well layers. *Phys. Rev. B* **95**, 144205 (2017).

50. Sharma, T. K. & Towe, E. On ternary nitride substrates for visible semiconductor light-emitters. *Appl. Phys. Lett.* **96**, 191105 (2010).

51. Michon, A. et al. Effects of pressure, temperature, and hydrogen during graphene growth on SiC(0001) using propane-hydrogen chemical vapor deposition. *J. Appl. Phys.* **113**, 203501 (2013).

52. Ben Jabra, Z. et al. Hydrogen-mediated CVD epitaxy of graphene on SiC: Implications for microelectronic applications. *ACS Appl. Nano Mater.* **4**, 4462 (2021).

53. Dussaigne, A. et al. Full InGaN red light emitting diodes. *J. Appl. Phys.* **128**, 135704 (2020).

## Acknowledgements

This work, carried out on the Platform for Nanocharacterisation (PFNC), was supported by the "Recherche Technologique de Base" and "France 2030 - ANR-22-PEEL-0014" programs of the French National Research Agency (ANR).

## Author contributions

A.D. wrote the manuscript, performed PL measurements and simulations. B.H., A.D. and B.D. conceived the project. C.P. performed the nitride growth experiments under the supervision of A.D and the graphene synthesis under the supervision of A.M. and S.V. C.P. performed XRD and Raman experiments. N.R. made CL measurements. D.C. performed HRTEM, HRSTEM, GPA, EDX, PED and electron holography on the red emitting TEM lamella. A.G. made HRSTEM and EDX experiments on the three different TEM lamella. A.D., B.H. and C.P. analysed the whole data. All authors contributed to the interpretation of the results. All authors revised the manuscript and approved the final version.

## Competing interests

The authors declare no competing interests.

## Additional information

**Supplementary information** The online version contains supplementary material available at https://doi.org/10.1038/s43246-024-00725-8.

**Correspondence** and requests for materials should be addressed to Amélie Dussaigne.

**Peer review information** *Communications materials* thanks the anonymous reviewers for their contribution to the peer review of this work. Primary Handling Editors: Jet-Sing Lee.

**Reprints and permissions information** is available at http://www.nature.com/reprints

**Publisher's note** Springer Nature remains neutral with regard to jurisdictional claims in published maps and institutional affiliations.

**Open Access** This article is licensed under a Creative Commons Attribution 4.0 International License, which permits use, sharing, adaptation, distribution and reproduction in any medium or format, as long as you give appropriate credit to the original author(s) and the source, provide a link to the Creative Commons licence, and indicate if changes were made. The images or other third party material in this article are included in the article's Creative Commons licence, unless indicated otherwise in a credit line to the material. If material is not included in the article's Creative Commons licence and your intended use is not permitted by statutory regulation or exceeds the permitted use, you will need to obtain permission directly from the copyright holder. To view a copy of this licence, visit http://creativecommons.org/licenses/by/4.0/.

© The Author(s) 2024