Optical properties of basal plane stacking faults in AlN nanowires

Corentin Guérin, Fabien Jourdan, Jean-Luc Rouvière, Bruno Gayral, Gwénolé Jacopin, Bruno Daudin

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Corentin Guérin, Fabien Jourdan, Jean-Luc Rouvière, Bruno Gayral, Gwénolé Jacopin, et al. Optical properties of basal plane stacking faults in AlN nanowires. Journal of Applied Physics, 2025, 137 (17), pp.174303. 10.1063/5.0252320. hal-05345072

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# Optical properties of basal plane stacking faults in AlN nanowires

Corentin Guérin 🐦 ; Fabien Jourdan; Jean-Luc Rouvière 🐦 ; Bruno Gayral 🐦 ; Gwénolé Jacopin 🐦 ; Bruno Daudin 🐦

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J. Appl. Phys. 137, 174303 (2025)

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# Optical properties of basal plane stacking faults in AlN nanowires

Cite as: J. Appl. Phys. **137**, 174303 (2025); doi: 10.1063/5.0252320

Submitted: 9 December 2024 · Accepted: 16 April 2025 ·

Published Online: 2 May 2025

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Corentin Guérin,<fnref n="1" />,<fnref n="2" />,ᵃ⁾ ID Fabien Jourdan,<fnref n="1" /> Jean-Luc Rouvière,<fnref n="3" /> ID Bruno Gayral,<fnref n="1" /> ID Gwénolé Jacopin,<fnref n="2" /> ID and Bruno Daudin<fnref n="1" /> ID

## AFFILIATIONS

## ABSTRACT

It is shown that depending on growth conditions, cathodoluminescence and photoluminescence spectra of AlN nanowires grown by plasma-assisted molecular beam epitaxy exhibit a peculiar optical signature around 5.85 eV. Several contributions to this peak can be identified, with a large variability in energy from nanowire to nanowire. We show that these peaks are correlated with the presence of I1 basal stacking faults (BSFs), crossing the entire diameter of nanowires. The formation of BSFs is assigned to the specific kinetics of Al and N adatoms when using migration enhanced epitaxy technique. By analogy with previous observations in GaN, we propose that I1 BSFs behave as zinc-blende quantum wells in a wurzite (WZ) matrix. Temperature and power dependence studies of the optical properties of these quantum wells by cathodoluminescence and photoluminescence are consistent with exciton coupling to randomly distributed donors in the WZ matrix and also suggest an inter-BSF coupling depending on their density.

© 2025 Author(s). All article content, except where otherwise noted, is licensed under a Creative Commons Attribution (CC BY) license (https://creativecommons.org/licenses/by/4.0/). https://doi.org/10.1063/5.0252320

## I. INTRODUCTION

III-N semiconductors, namely, GaN, AlN, InN, and their alloys, are characterized by zinc-blende/wurtzite polytypism.<fnref n="1" /> Although the most stable wurtzite (WZ, hexagonal) phase consists of an ABABAB... stacking sequence along the (0001) direction, the zinc-blende (ZB, cubic) structure consists of an ABCABC... stacking sequence along the (111) cubic direction. The higher symmetry of the cubic variant, which leads to an absence of spontaneous polarization, and the higher electron mobility in n-type ZB GaN<fnref n="2" /> have motivated for long a continuous interest in ZB GaN and InGaN in view of realizing high efficiency optoelectronic devices. In particular, as the bandgap value of the ZB variant is smaller than the bandgap value of the WZ phase,<fnref n="3" /> stacking faults (SFs), which can be viewed as ultrathin ZB GaN insertions in the WZ GaN matrix, were soon identified as binary quantum wells emitting in the UV range.⁴,⁵ Unfortunately, the potential of such cubic III-nitride quantum wells (QWs) appears difficult to concretize due to the difficulty to control at the monolayer (ML) scale the switching from WZ to ZB and vice versa.

In practice, ZB GaN is grown at temperatures lower than its WZ counterpart,⁶⁻⁸ which further complicates the precise tuning of the thickness/emission energy of ZB insertions in the WZ matrix.

However, the recent interest in III-nitride nanowires (NWs) has put in evidence an interesting peculiarity of these 3D nano-structures with respect to layers. Indeed, the growth of InGaN and GaN NWs under metal-rich conditions leads to their widening, which is associated with SFs' formation.⁹⁻¹¹ Then, in NWs, the controlled formation of SFs can in principle be advantageously achieved by tuning the metal/active N ratio value, while maintaining a constant growth temperature over the whole growth process, which should facilitate the controlled formation of ZB GaN QWs. Nevertheless, as in layers, several types of SFs have been identified in GaN NWs, namely, I1, I2, and E (extrinsic) SFs depending on the number of planes deviating from the WZ stacking sequence.¹²,¹³ The relatively low formation energy of such SFs¹²,¹³ is consistent with their simultaneous observation in GaN NWs, which was assessed by a combination of high-resolution electron microscopy and optical measurements.¹⁴⁻¹⁶ On the one hand, in NWs that

<footnote n="1">¹Univ. Grenoble Alpes, CEA, Grenoble INP, IRIG, PHELIQS, NPSC, 17 av. des Martyrs, Grenoble 38000, France</footnote>

<footnote n="2">²Univ. Grenoble Alpes, CNRS, Grenoble INP, Institut Néel, Grenoble 38000, France</footnote>

<footnote n="3">³Univ. Grenoble Alpes, CEA, Grenoble INP, IRIG, MEM, LEMMA, 17 av. des Martyrs, Grenoble 38000, France</footnote>

<footnote>ᵃ) Author to whom correspondence should be addressed: corentin.guerin@cea.fr</footnote>
are free of extended defects such as threading dislocations or grain boundaries, SFs are forming "perfect imperfections" quantum wells crossing the full NW section.¹⁷⁻²⁰ On the other hand, the simultaneous observation of SFs or cubic section exhibiting different thicknesses again prevents from finely tuning their optical properties in a controlled way.

The situation is different in AlN where the high formation energy of SFs makes more favorable the formation of I1 SFs,¹³ consistent with experimental observations in AlN layers.²¹ One motivation of the present work is to investigate the optical properties of such SFs in AlN NWs, a field still virgin in the literature. It will be shown that SFs can be produced in a controllable way in AlN NWs, depending on growth conditions, opening the path to the realization of binary ZB AlN QWs in a WZ AlN matrix for the development of a new type of far UV-emitting devices.

## II. EXPERIMENTAL DETAILS

The common substrate for our samples consists of patterned GaN NWs, grown by metalorganic chemical vapor deposition (MOCVD) on an Si substrate. These GaN stems are 300 nm in diameter for 1 µm in height. After introduction in the growth chamber and outgassing at 350 °C during 12 h, AlN NW sections were grown by plasma-assisted molecular beam epitaxy (PA-MBE) directly on the top of GaN stems. The Al and N flux values were equivalent to 0.6 Ml/s, as determined by RHEED oscillations on a layer of AlN grown in Al-limited and N-limited conditions, respectively. Sample A is a control sample. It was grown under standard conditions, i.e., at a substrate temperature of 950 °C and an Al/active N ratio of 1 corresponding to a growth rate of 0.6 Ml/s during 90 min. The second sample (sample B) was grown by migration enhanced epitaxy (MEE) technique, i.e., by alternating Al and N flux (20 s/30 s) for 280 cycles. In this case, Al is limiting the growth rate. Therefore, assuming negligible Al desorption from the surface, 280 cycles correspond to 93 min of growth (total time with the Al cell open). As previously demonstrated, this technique has a radical effect on the morphology quality and optical properties of AlN NWs.²² An additional sample (sample C) was designed to be a combination of samples A and B: An AlN section was grown for 90 min using the same conditions as for sample A followed by the deposition of a 20 nm GaN spacer and deposition next to the AlN section grown using the same conditions as for sample B.

Cathodoluminescence (CL) spectroscopy was carried out on samples A and B using a modified scanning electron microscope (SEM). In this system, the emitted light is collected by a parabolic mirror and refocused to the entrance slit of a 550 nm focal length spectrometer equipped with a 1800 grooves/mm diffraction grating. The luminescence signal is then recorded by a UV-enhanced CCD camera at the output of the spectrometer. All measurements were performed at 5 K with an acceleration voltage of 5 kV in order to limit the excitation volume to the AlN NW sections. The beam current was approximately 1 nA. The entrance slit of the spectrometer was set to 0.1 mm, resulting in a spectral resolution of around 0.1 nm. Unless specified otherwise, each spectrum was acquired by focusing the electron beam on the top of a single NW.

Scanning transmission electron microscopy (STEM) observations were performed on a Thermofisher TITAN-Themis microscope

having a Cs probe-corrector. The convergence angle was set to 24 mrad in order to have high-resolution images along the a = [2-1-10] zone axis. 8 K × 8 K high-resolution scanning transmission electron microscope (HR-STEM) images were acquired, thanks to the Gatan scanning unit and Digital Micrograph software. These 8 K × 8 K allowed us to use low magnification modes (from ×300 to ×900 k) in order to have both a large field of view and high-resolution (HR) images showing Al and N atomic columns.

Photoluminescence (PL) spectra were acquired using a 195 nm laser light source from a pulsed quadrupled titanium-sapphire laser. The pulse duration is about 150 fs, and the pulse repetition rate is 80 MHz. The luminescence signal is collected by a fused silica lens and then focused by another fused silica lens to the entrance slit of a 460 mm focal length spectrometer equipped with a 1800 grooves/mm diffraction grating. The luminescence signal is detected by a liquid nitrogen cooled UV-enhanced silicon CCD.

## III. RESULTS AND DISCUSSION

### A. Origin of the 5.85 eV signal

First, CL has been performed on samples A and B. We can observe in Fig. 1(a) that near band edge (NBE) emission energies of both A and B samples are similar. The two intense peaks, separated in both cases by about 30 meV, correspond to free exciton (FE) recombination ($I_{5}^{n=1}$) and recombination of donor-bound exciton ($D_{0}X_{1}$), already observed in the literature.²³ Noticeably the absolute values of the peaks are slightly different in the two probed wires, with the free exciton peak for sample A (B) observed at 6.040 eV (6.045 eV) and the $D_{0}X_{1}$ at 6.010 eV (6.015 eV), which is assigned to a different stress state. For sample A (in black), we observe weak peaks at 110 and 230 meV below the FE recombination peak, which are attributed to phonon replica, respectively, 1LO and 2LO, of the FE recombination peak.

Interestingly, for sample B (in red), we observe an intense doublet at 5.85 eV. This doublet cannot be attributed to a phonon replica of the NBE structure, as it has a distinctly different spectral shape from the NBE. Moreover, this peak is itself followed by two replica at 110 and 225 meV. These replica share the same “doublet” shape as the peak at 5.85 eV and are, therefore, assigned to 1LO and 2LO phonon replica of this peak.

STEM observations have been conducted on focused ion beam (FIB)-prepared cross section lamellas of both samples, with a focus on detecting any defect specific to sample B. The high-resolution STEM (HR-STEM) images were analyzed using the geometrical phase analysis (GPA) approach.²⁴ It consists of applying a mask on the fast Fourier transform of the HR-STEM image to filter a single family of crystallographic planes and calculate the corresponding amplitude [shown in Fig. 1(b)] and phase (not shown here) images. In Fig. 1(b), we observed three planar defects in MEE-grown AlN section of sample B, indicated by discontinuities in the amplitude image. These defects were consistently found in the majority of the examined wires in sample B and always extended across the entire diameter. In contrast, no such defects were detected in the AlN section of sample A [see Fig. 1(c)].

Based on the HR-STEM image shown in Fig. 1(d), we conclude that these defects are I1 basal stacking faults (BSFs) corresponding to one unique faulted plane along the c-plane stacking
<figure><img src="image_11.png" /><figcaption>FIG. 1. (a) Cathodoluminescence spectra of standard MBE (black)- and MEE (red)-grown AlN nanowires. (b) Geometric phase analysis (GPA) of a cross-sectional HR-STEM image of a MEE-grown NW. (c) GPA amplitude of a cross-sectional HR-STEM image corresponding to standard MBE-grown NW. (d) HR-STEM image of a 2D defect in MEE-grown sample B, corresponding to the defective region and allowing one to identify an $I_1$ basal stacking fault (BSF) of ABABCBCB type.</figcaption></figure>

sequence: ABABABCBCBC. Considering that the 5.85 eV peaks just appear for sample B, it suggests that they are related to I1 BSF observed in Figs. 1(b) and 1(d). The fact that none of them are observed on sample A [Fig. 1(c)] reinforces this hypothesis.

No other type of fault has been observed in the samples under scrutiny, consistent with Stampfl and Van de Walle's calculation on the formation energy of stacking faults in AlN.¹³ Specifically, their results indicate that the formation energy of I1 BSF in AlN (49 meV) is two times smaller than of I2 (103 meV) and E (97 meV) BSF. Therefore, I1 BSFs are much more likely to occur in AlN than other types of BSFs, as observed by Dovidenko et al. in the case of layers.²¹

Interestingly, the presence of BSFs is correlated to MEE growth conditions, pointing out the role of kinetics. Adatom

kinetics on Al-polar (0001) AlN has been investigated by Jindal and Shahedipour-Sandvik,²⁵ under Al-rich conditions, (0001) surface diffusion barriers for both Al and N are found to be lower than on the N-rich surface. Due to the high growth temperature, far above the Al melting point, Al mobility is indeed expected to be dominated by intermetallic interaction and then very large compared to N, under the extreme Al-rich conditions corresponding to the Al-exposure step of the MEE cycle. This favors the occupancy of the low energy hcp sites and WZ growth when shifting to the N-exposure step. However, at the end of this step, the increasing N coverage simultaneously with the reduction in Al coverage leads to a considerable increase in both Al and N surface diffusion barriers. Then, the probability that Al adatoms trapped in the fcc position bound to N to form a ZB nucleus is expected to increase.
This mechanism should be favored by the similar energy of fcc and hcp sites in the case of an N-rich surface, making both sites equally favorable to Al adsorption. As further shown recently by Zang et al.,²⁶ the stability of WZ and ZB nuclei depends on their size, eventually becoming favorable for the ZB phase for the smallest size. We speculate that SFs' formation when using MEE growth conditions should preferentially occur at one moment of the cycle corresponding to the rapid transition from Al-rich to N-rich conditions associated with an abrupt change in the Al adatom surface diffusion barrier.

To confirm the link between I1 BSF and the unknown peak of Fig. 1(a), we performed HR-STEM observations and CL on a focused ion beam (FIB)-prepared lamella of sample C. As already said, sample C consists of a mixture of both standard and MEE growth modes. As expected, BSFs are localized in the top part of this sample [see Fig. 2(b)], while no BSFs are observed in the lower part (not shown).

The CL vertical linescan, derived from hyperspectral mapping²⁷ and integrated across 30 points along the horizontal black dashed double arrow in Fig. 2(a), reveals that the BSF signal originates from

<figure><img src="image_12.png" /><figcaption>FIG. 2. (a) Schematic and SEM image of an FIB-prepared lamella from sample C. (b) Zoom of the GPA amplitude of the HR-STEM image of the top AlN of the wire, high-lighting the presence of several BSFs. The faint vertical dark lines in the picture are a scanning artifact. (c) CL linescan of the probed NW, emphasizing the correlation between SFs and the peak at 5.88 eV (for this wire, the peak was unique, not a doublet, and centered at this energy). The white dashed line shows the position along the growth axis. The scanline shown here is the integration of the signal from 30 points laterally scanned along the black dashed double arrow shown in (a).</figcaption></figure>
the top of the MEE-grown AlN section. This observation strongly supports its attribution to I1 BSFs. Noticeably, we observe AlN NBE sharp line all along the wire, assigned to the presence of an AlN shell. Moreover, we see that NBE is particularly intense, with a stronger $D_0$X1 contribution in MEE AlN, compared to the conventionally grown lower AlN section. As previously observed,²² this confirms that MEE drastically improves AlN structural and optical properties. Interestingly, one furthermore observes an anti-correlation between the $D_0$X1 signal and the one assigned to BSFs, in the MEE-grown top AlN section, indicating that carriers are efficiently collected by BSFs. This suggests that ZB I1 BSFs act as thin QWs in a WZ matrix where the recombination of bulk AlN excitons is occurring.

## B. Temperature dependence of the I1 BSF signal

In order to better understand the charge carrier trapping mechanisms by BSFs, we then performed temperature-dependent CL studies on sample B. In Fig. 3(a), we report the variation in the CL spectrum in a spectral region centered at 5.85 eV as a function of temperature. This study was done on a 20 × 20 µm zone, which corresponds to about 625 NWs. Then, the resulting CL spectrum can be viewed as an average response, integrating CL variability from NW to NW. This variability of the NW optical response will be discussed later.

First, it appears that the peak we identified as $D_0$X1 dominates at 300 K to the expense of FE contribution, which is progressively vanishing. Nevertheless, the 30 meV energy difference between the

two peaks at 5 K is in accordance with the 28 meV reported by van Deurzen et al.²³

The temperature dependence of the 5.85 eV BSF peak energy position is plotted in Fig. 3(b) alongside the energy position of the other NBE contributions. Interestingly, we observe a difference of around 150 meV between the FE emission peak and the BSF-related luminescence peak at 100 K. This energy corresponds to the binding energy of the FE to the BSF QW, significantly larger than for GaN (45 meV from Rebane²⁸). Belabbes et al.²⁹ have proposed that ZB AlN inclusions inside the WZ matrix are type II heterojunctions. They calculated a difference between conduction bands of $\Delta E_c = 1.181$ eV between the X ZB conduction band minimum and the Γ WZ conduction band minimum. Thus, the electron will be confined in the ZB AlN inclusion at X, while the hole will be delocalized in the WZ AlN at Γ. The spatially indirect character of the transition is consistent with the presence of strong phonon replica.

Despite the spatial separation induced by the type II character, the hole remains in the vicinity of the ZB QW due to the combined action of Coulomb interaction (hole attraction by the electron bound to the BSF) and the strong internal electric field induced by the spontaneous polarization mismatch between WZ and ZB AlN.

In line with this interpretation, the BSF-related emission rapidly decreases with increasing temperature, suggesting a weak confinement of the electron within the thin ZB inclusion. Once the electrons gain enough thermal energy to escape the ZB inclusion, the BSF emission disappears. Experimentally, this signal vanishes around 150 K. In Fig. 3(b), the superposition of FE luminescence data shifted by 150 meV on BSF peak data points toward a slight

<figure><img src="image_13.png" /><figcaption>FIG. 3. (a) Energy emission of BSF-related CL peak from 5 to 300 K for a 20 × 20 µm scanning zone (≈625 wires). (b) Variation in the peak energy of FE (blue), $D_0$X1 (black), BSF peak (green), and BSF peak first phonon replica (red) with CL excitation. The blue line corresponds to the free exciton energy down-shifted by 150 meV.</figcaption></figure>
<figure><img src="image_14.png" /><figcaption>FIG. 4. Variability of I1 peak shape between different wires at 5 K in sample B. NWs were probed one by one with 5 kV excitation focused on the c-plane top facet.</figcaption></figure>

localization effect at low temperature. This corresponds to an esti-
mated electron localization energy in the BSF of approximately
10 meV.

C. Exciton localization effects

As further shown in Fig. 4 in confirmation of data reported in
Figs. 1 and 3, BSF-related peaks present a marked variability from
NW to NW. Most spectra exhibit a doublet or, less frequently, a
triplet and/or shoulders suggesting the presence of additional
peaks. However, the variability of the peak wavelength emission

shown in Fig. 4 as well as the possible exciton localization put in
evidence by the temperature dependence of the BSF-related peak
shown in Fig. 3(b) cannot be assigned to the presence of steps or to
thickness fluctuations because of the “perfect” 2D structure of AlN
BSFs extending through the entire diameter of the wire, with a
sharp ZB/WZ interface. This leads to questions about the origin of
variability in the emission and exciton localization.

In agreement with the model proposed by Corfdir et al.,³⁰ we suggest that the trend to localization revealed by the dependence of the BSF-related emission peak position in Fig. 3(b) is correlated with electron localization in the BSF due to the nearby presence of donors randomly distributed in the surrounding WZ matrix.

Assuming a similar mechanism in the case of AlN BSFs, it is
tentatively suggested that O, which is a well identified residual
donor in AlN³¹ with a luminescence signature around 3.8 eV,²²,³²,³³
could be responsible for the electron localization in ZB QWs under
investigation. As a matter of fact, the MEE step corresponding to
exposure to Al is particularly favorable to residual O trapping at
levels higher than for standard growth conditions.²² As shown in
Fig. 2(b), the average spacing between I1 BSFs in AlN NWs lies in
the 5–20 nm range, small enough to ensure that most O impurities
in this case should be in the close vicinity of individual BSFs and
could contribute to electron localization.

A variability in the width and shape of the BSF-related peaks
seen in Fig. 4 has been similarly observed by Corfdir et al. in
GaN.³⁴ It was demonstrated that a significant coupling between
BSFs can occur until 10 nm, resulting in variations in the effective
QW depth and barriers.³⁵ In fact, as shown in Fig. 2, we observe
BSF bundles, with a separation between adjacent BSFs lower than
10 nm, supporting the hypothesis of inter-BSFs coupling. Such a
coupling has also been proposed by Paskov et al.³⁶ and by
Lähnemann et al.¹⁶ for GaN BSFs. The order of magnitude
(between 30 and 50 meV) they observed between peaks assigned to
BSF bunching is similar to what we found in the present study.

<figure><img src="image_15.png" /><figcaption>FIG. 5. (a) Evolution of PL spectra with temperature. The spot size is about 30 µm, corresponding to about 1400 NWs. (b) Evolution of the I1 BSF peak with the injection power density at 5 K.</figcaption></figure>
Along the same line, we propose that the presence of a large number of BSFs, arranged in diverse configurations, in one AlN NW, could be the cause of this variability in the I1 BSF peak emissions. Finally, it cannot be excluded that both inter-BSF and donor-BSF coupling could occur here and explain the variability in the peaks shown in Fig. 4. Further studies should be necessary to de-correlate the effects of both contributions.

### D. Power dependent photoluminescence

Due to the polar character of the AlN WZ phase, it is expected that the optical properties of I1 BSF should be affected by the presence of an internal electric field. In order to clarify this point, PL measurements as a function of both temperature and excitation power were performed.

First, in Fig. 5(a), it appears that the maximum of the BSF-related PL emission does not shift with temperature until about 100 K, as already observed in Fig. 3. The variation in the BSF-related PL emission peak as a function of laser power is further shown in Fig. 5(b) and reveals a slight blue shift for increasing injection power.

This blue shift is assigned to the screening of the electric field at the interface of the ZB quantum well included in the WZ matrix responsible for the so-called quantum confined Stark effect (QCSE). It is directly related to the difference in surface charge density between the polar surface of the WZ AlN and the non-polar (111) ZB AlN. It has to be noted that this difference in polarization is three times larger in AlN than in GaN, with the consequence of being non-negligible even for 3-ML thick QWs. As a direct consequence of the type II QW nature of ZB AlN insertions in WZ AlN theoretically predicted by Belabbes et al.,²⁹ the delocalization of the holes in the WZ barrier should cause the apparition of a dipole in the QW when carriers are injected into the structure. This dipole tends to screen the field induced by QCSE, causing emission blue shift. Such a screening has already been observed by Lähnemann et al.¹⁶ and Jacopin et al.¹⁴ by changing the distance (and thus the amount of injected carriers) between the CL spot and the BSF position in GaN.

## IV. CONCLUSION

In summary, we have shown that the specific adatom kinetics induced by MEE growth conditions favors the formation of I1 BSFs. These BSFs, which correspond to a single faulted plane in the WZ stacking sequence, lead to the homoepitaxial formation of ultrathin ZB AlN QWs through the full diameter section of AlN NWs. By correlating HR-STEM observations and CL experiments in single NWs, we have found that the emission band around 5.85 eV is related to the presence of these I1 BSFs. Such an assessment was further supported by power dependent PL measurements exhibiting an emission blue shift for increasing power, assigned to screening of the internal electric field generated by the polarization difference at the ZB/WZ interface. Similar to the widely studied case of GaN, the low temperature carrier localization regime revealed by temperature-dependent CL and PL was assigned to the coupling of the exciton in the QW with residual donors randomly distributed in the surrounding WZ matrix. This hypothesis was supported by the chromatic variability of the 5.85 eV emission band from NW to NW, eventually modulated by inter-BSF

coupling depending on their density and average spacing. As BSF formation can be triggered by the control of the flux alternance cycle in MEE, these results also pave the way to controlled emission in the UV-C range of ZB/WZ homoepitaxial AlN NW heterostructures, despite the quenching of SF-related emission at 160 K, which requires additional investigation to be overcome.

## ACKNOWLEDGMENTS

The authors acknowledge the support from the French ANR agency (Project ANABASE, No. ANR-22-CE09-0031). They also acknowledge the support from GANEXT (No. ANR-11-LABX-0014). GANEXT belongs to the public funded "Investissements d'Avenir" program managed by the French ANR agency. The authors thank Fabrice Donatini for his help with cathodoluminescence measurements.

## AUTHOR DECLARATIONS

### Conflict of Interest

The authors have no conflicts to disclose.

### Author Contributions

**Corentin Guérin:** Conceptualization (lead); Data curation (lead); Formal analysis (lead); Investigation (lead); Methodology (equal); Software (lead); Visualization (lead); Writing – original draft (lead); Writing – review & editing (equal). **Fabien Jourdan:** Data curation (equal); Formal analysis (equal); Methodology (equal). **Jean-Luc Rouvière:** Data curation (equal); Software (equal); Visualization (equal). **Bruno Gayral:** Data curation (equal); Software (equal); Visualization (equal). **Gwénolé Jacopin:** Conceptualization (equal); Funding acquisition (equal); Methodology (equal); Resources (equal); Supervision (equal); Writing – review & editing (equal). **Bruno Daudin:** Conceptualization (equal); Funding acquisition (lead); Methodology (equal); Project administration (lead); Resources (lead); Supervision (lead); Writing – original draft (equal); Writing – review & editing (lead).

## DATA AVAILABILITY

The data that support the findings of this study are available within the article.

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