# Assessment of the crystalline orientation of nanoscale semiconductor structures via atom probe tomography

J. Cañas, A. Grenier, J L Rouviere, A. Harikumar, S. Ndiaye, A. Jannaud, E.

Monroy, Lorenzo Rigutti

► To cite this version:

J. Cañas, A. Grenier, J L Rouviere, A. Harikumar, S. Ndiaye, et al., Assessment of the crystalline orientation of nanoscale semiconductor structures via atom probe tomography. *Applied Physics Letters*, 2025, 126 (11), pp.112101. 10.1063/5.0242659. *hal-04994349*

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Assessment of the crystalline orientation of nanoscale semiconductor structures via atom probe tomography

J. Cañas¹,ª), A. Grenier², J.L. Rouviere³, A. Harikumar¹, S. Ndiaye⁴, A. Jannaud², E. Monroy¹ and L. Rigutti⁴

¹Univ. Grenoble-Alpes, CEA, Grenoble INP, IRIG, PHELIQS, 38000 Grenoble, France

²Univ. Grenoble Alpes, CEA, LETI, 38000 Grenoble, France

³Univ. Grenoble Alpes, CEA, IRIG, MEM, 38000 Grenoble, France

⁴UNIROUEN, CNRS, Groupe de Physique des Matériaux, Normandie Université, 76000 Rouen, France

ª) Author to whom correspondence should be addressed: jesus.canas@neel.cns.fr

Abstract

We demonstrate the application of atom probe tomography for assessing the crystalline orientation of nanoscale semiconductor structures via the analysis of charge state ratio maps in the detector space. The experimental realization is carried out in the context of adventitious cone-shaped domains present in AlGaN quantum dot superlattices. The cone-shaped domains, which emerge from shallow pits generated in AlN and propagate through the superlattices, are shown to exhibit small misorientation angles of their crystalline <0001> poles. The results of the atom probe tomography analysis are confirmed by convergent beam electron diffraction measurements. The use of this methodology adds another layer to the application of this technique to semiconductor nanoscale systems, providing not only compositional maps, but also information on the crystallographic orientation.

Atom probe tomography (APT) is a technique enabling the three-dimensional (3D) compositional mapping with sub-nanometer resolution of a nanoscale volume. It is based on the controlled field evaporation of the surface atoms of a sharp (apex diameter ≲ 100 nm) tip, and their analysis through a position- and time-of-flight-sensitive detector. The information collected during the field evaporation of the specimen is subsequently analyzed to reconstruct the 3D distribution of chemical species. The reconstruction allows accessing nanoscale sizes, shapes, and chemical composition of 3D features within the field evaporated volume.¹ Laser-assisted atom probe tomography (La-APT) extended the capabilities of APT enabling the study of non-conductive specimens, including semiconductors.² La-APT can be used to study semiconductor nanostructures such as quantum well or quantum dots (QDs),³ which play an important role in the development of efficient photon emitters (lamps, lasers or single photon sources) from the far-UV to the NIR. In the case of QDs, as nanoscale 3D objects, their study with transmission electron microscopy (TEM) or scanning tunneling microscopy (STM) is incomplete as it is limited to a 2D projection or slice, respectively. On the contrary, La-APT can provide 3D structural and compositional information, which is fundamental to understand their optoelectronic properties. Therefore, this technique can serve as a powerful tool for addressing
semiconductor challenges such as the efficiency drop observed in far-UV or green/amber emitting devices. La-APT is also well suited to study structural extended defects or nanoscale domains in semiconductors, especially those that appear associated with compositional fluctuations, such as grains, pits, or grooves at heterointerfaces. For instance, in our earlier contribution, we utilized La-APT to probe the structure and compositional variations of cone-shaped domains existing in AlGaN/AlN QDs superlattices (SLs), which were associated with bimodal emission.⁴ Furthermore, La-APT has been applied to the study of alloy composition and impurity segregation at dislocations⁵,⁶ or stacking faults⁷,⁸. However, there is little awareness in the semiconductor community that APT can probe the crystalline orientation of nanostructures in 3D.⁹,¹⁰ In metallurgy, several low-index poles and zone lines can often be identified from desorption maps¹¹. This is generally not the case in semiconductors¹²,¹³, where the spatial resolution is lower. However, the measurement of crystal orientation of nanostructures can be performed by identifying the main low-index pole through the analysis of selected ion density patterns, ion or atomic fractions, or charge state ratios (CSR). In the case of III-N materials, it has been shown that CSR(Al) or CSR(Ga) patterns in the detector space can indeed help identify the position of the <0001> pole and thus its orientation with respect to the specimen axis³,¹⁴,¹⁵. This is because the CSR is correlated to the evaporation field, which is enhanced at the edges of atomic terraces formed around low-index poles. Therefore, the CSR maps in the detector space provide a representation of the electric field on the tip surface, which can be used to locate the <0001> pole in III-N materials. Yet, the distribution of the CSR in the detector space may exhibit diverse spatial dependence, related to the experimental conditions and to the quality of the specimen. Distributions with nearly circular symmetry³,¹⁴, as well as distributions exhibiting a well-resolved six-fold symmetry, as expected in hexagonal crystal structures¹⁶, have been reported. Furthermore, the symmetry of the CSR patterns of the <0001> orientation is significantly different from those of non-polar directions.¹⁷

In this letter, we use atom probe tomography (APT) to observe the misorientation of cone-shaped domains in AlGaN/AlN QD SLs designed for UV emission. The crystallographic analysis conducted by APT is validated by convergent beam electron diffraction patterns obtained by transmission electron microscopy. This methodology establishes the utilization of APT for investigating crystalline (mis)orientation alongside the compositional and structural characteristics of semiconductor nanostructures.

The studied sample consists of 100 periods of Al<sub>x</sub>Ga<sub>x-1</sub>N/AlN self-assembled QDs deposited on commercial 1-µm-thick (0001)-oriented AlN-on-sapphire templates by plasma-assisted molecular beam epitaxy (MBE). For the QD epitaxy, the metal/N flux ratio was kept below the stoichiometric value, resulting in the Stranski-Krastanov (SK) growth mode.<sup>18</sup> QDs with a diameter of 6.2 ± 1.2 nm and a height of 1.3 ± 0.2 nm were formed by deposition of 5 ML of Al<sub>0.14</sub>Ga<sub>0.86</sub>N. The AlN overgrowth to form
the barriers was performed under Al-rich conditions (Al/N ≈ 1.1) followed by a growth interruption under nitrogen, leading to a total barrier thickness of ≈ 4.4 nm. The aluminum-rich conditions were determined based on flux calibrations performed within the MBE system. The grown AlN layers are always stoichiometric, and any excess metal accumulates on the growing surface. RHEED oscillations are damped with metal-rich conditions, so the Al/N ratio can be measured by comparing the time required to consume the excess of metal with the layer growth time. The sample presents bimodal emission with a main emission wavelength at 291 nm and a secondary emission at 302 nm caused by a secondary population of QDs which are associated with cone-shaped domains emerging from the AlN layers. Further information on the synthesis, structure and emission properties of the sample can be found in our previous works.⁴,¹⁹ Additionally, the sample structure and the cone-shaped domain characteristics are detailed in the **supplementary material**. A TEM lamella and a needle-shaped specimen for APT were prepared using a scanning electron microscope/focused ion beam (SEM/FIB) Helios Nanolab 450S operated at 16 kV. Low energy (2 keV) ion beam was used during the final APT annular milling to reduce irradiation damage on the sidewalls of the specimen.

The structural properties of the samples were studied by bright field (BF) and high-angle annular dark field (HAADF) scanning transmission electron microscopy (STEM) imaging using a probe-corrected Titan Themis (Thermo Fisher Scientific) operating at 200 kV. Convergent beam electron diffraction (CBED) discs were recorded on the <1 1 -2 0> zone axis of samples within the cone-shaped domains and homogeneous regions to access the crystalline orientation of the material. In CBED patterns, small variations in crystalline orientation can be precisely measured by analyzing the displacement of the center of the Ewald sphere. The reciprocal lattice points that lie on the Ewald sphere correspond to crystal lattice planes that satisfy the Laue condition for diffraction. Only diffraction spots from the zero-order Laue zone (ZOLZ)—that is, only lattice planes aligned with the incident electron beam—contribute significantly to the central part of the CBED pattern. As the sample orientation or crystallography changes, different sets of reciprocal lattice points will coincide with the Ewald sphere. Consequently, the intersection of the Ewald sphere with the ZOLZ, the ZOLZ Ewald circle (ZEC), will shift in the CBED patterns with changes in crystallographic orientation. Figure 1a presents a schematic representation of the displacement of the ZEC in the reciprocal space upon the tilt of the observed electron-transparent specimen.

A needle-sharp specimen containing a cone-shaped domain was analyzed by La-APT using a CAMECA FlexTAP system equipped with a pulsed laser operating at 343 nm, with a pulse duration of 450 fs, an energy of 16 nJ/pulse and a repetition rate of 100 kHz. The laser has a diameter of roughly 60 µm at the field evaporated tip, corresponding to a moderate energy density of approximately 0.5 mJ/cm². The tip temperature was maintained around 40 K during analysis. The field of view was ±15°;
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roughly equivalent to a straight flight atom probe with L=74 cm flight-path. Previous studies have suggested that for AlGaN and AlN, the measured composition either does not depend on the laser pulse energy or shows only a very weak dependence.3,15 This suggests that the AlGaN system is not prone to thermally driven inhomogeneous decomposition during field evaporation. A common feature, however, is that the electric field strongly influences the measured composition. The experiment was performed under electric field conditions that mitigate these effects, as suggested by the charge-to-state ratio $Ga^{2+}/Ga^{+} \approx 0.03$ (the mass/charge spectrum is shown in the supplementary material), corresponding to an average field of 23 V/nm at the tip's surface. This protocol results in a high signal-to-noise ratio in the mass/charge spectrum and reduces the preferential loss of metallic species during the tip evaporation9,20. APT data were reconstructed using a standard cone angle reconstruction algorithm with a 12° cone-angle and 21 nm tip curvature radius. This methodology enables the transformation of the detector space positions (X, Y, N), where N is the ion sequence number, into the reconstructed coordinates of the specimen ($x_s, y_s, z_s$). The field-evaporated specimen is modeled as a spherical surface on a truncated cone, with ion ($x_s, y_s$) coordinates computed using point-projection and a magnification factor, and the $z_s$ coordinate adjusted incrementally for each ion considering a detector efficiency of 25%.

In the schematic representation of an La-APT experiment in figure 1b, the atoms of the specimen tip that protrude the most from the surface experience the strongest electric field, making them the first to field evaporate. The surface roughness associated with poles, i.e., low-index crystallographic planes forming annular terraces around the corresponding crystallographic directions, enhances the electric field and induces inhomogeneities in its distribution. In semiconductors, it is in general possible to identify one main low-index pole through the analysis of selected ion density patterns, ion or atomic fractions, or CSR. The CSR(Me) of a given metallic element Me is defined as:

$$CSR(Me) = \frac{n(Me^{k+})}{n(Me^{(k-1)+})}, \qquad (1)$$
Where Me<sup>k+</sup> corresponds to the k-th ionization state of the element Me, and n is the number of detection events within a user-defined volume. The CSR is an indicator of the microscopic surface field, as described in the framework of the post-ionization theory.<sup>21</sup> Furthermore, unlike composition-based maps, CSR maps exhibit a pattern with an approximately composition-independent symmetry, which makes it particularly interesting in specimens with composition inhomogeneities, as it is the case in tips containing multiple heterostructures. Therefore, CSR detector space maps can be used to identify the primary low-index pole, since it corresponds to the region where electric field is maximum, along the evaporated tip volume, providing an estimate of the crystallographic orientation, as illustrated in figure 1b.

<figure><img src="image_3.png" /><figcaption>FIG. 1. a) Schematic representation of the displacement of the zero-order Laue zone (ZOLZ) Ewald circle (ZEC) in the reciprocal space due to the rotation of the electron-transparent thin specimen in the direct space. b) Atom probe tomography schematic representation. A sharp needle tip under a high V<sub>DC</sub> field-evaporates atoms responding to a laser pulse. The main low-index pole (<0001> pole in the case of polar wurtzite nitrides) displays a local increase in the electric field that can be assessed though the charge-to-state ratio in the detector space maps.</figcaption></figure>

Figure 2a presents a STEM-HAADF image of the base of the AlGaN/Al QD SL, including the top of an AlN buffer layer beneath. The QD planes, which appear brighter in the image due to their higher Ga composition, resemble quantum wells because of the projection of several small QDs and their wetting layers throughout the depth of the lamella (approximately 50 nm). A cone-shaped domain is visible in the center of the image. It is characterized by QD planes that are bent and shifted relative to the rest of the sample. This domain originates from a shallow pit with a flat base and inclined, faceted walls in the AlN layer. This shallow pit morphology propagates as the SL grows and generates cone-shaped domains with a core where the QD planes are horizontal but vertically shifted compared to the rest of the sample, and a boundary region where the QD planes are bent to follow the inclined, faceted walls of the pit. These domains usually have a cone-like shape, as their lateral dimension tends to decrease along the growth axis. The structure of the cone-shaped domains was the focus of our previous work.<sup>4</sup>
PLEASE CITE THIS ARTICLE AS DOI: 10.1063/5.0242659

CBED patterns were acquired on the <11-20> zone axis in the AlN layers for both the homogeneous region and the cone-shaped domain, as represented in figures 2b and 2c, respectively. These CBED patterns correspond to wurtzite crystallographic structure, with the [0001] growth direction depicted vertically and the [0110] direction depicted horizontally. At the center of the ZEC, the superposition of the transmitted beam and the constructive interference result in maximum intensity in the diffraction spots. In the displayed CBED, the center of the ZEC has been determined as the center of mass of the diffracted spots' intensity, indicated with a dotted cross within the equi-intensity circumference. Upon comparing figures 2b and 2c, a clear displacement of the center of the ZEC is observed. This displacement arises within the cone-shaped domain due to the tilt of the <0001> axis with respect to its position in the homogeneous areas of the sample. However, it is difficult to determine whether the crystalline misorientation originates from the core or the boundary of the cone-shaped domain because the convergent electron beam traverses both regions in its path through the lamella.

<figure><img src="image_4.png" /><figcaption>FIG. 2. (a) STEM-HAADF image showing a cone-shaped domain propagating from the AlN buffer layer (at the bottom) into the AlGaN/AlN SL (at the top). Convergent beam electron diffraction (CBED) patterns obtained by localizing the beam (b) in an AlN barrier outside of the cone-shaped domain, and (c) within the cone-shaped domain. The respective locations are marked as cyan spots in (a). The center of the ZEC is indicated by the dotted cross within the equi-intensity circumference.</figcaption></figure>

The reciprocal space vector joining the two ZEC centers is $\bar{D} = (0.37 \AA^{-1}, 0.20 \AA^{-1})$ with a norm $||\bar{D}|| = 0.42 \AA^{-1}$. The angle between the two crystalline orientations can be calculated by applying the Bragg law for small angles $\theta = \lambda_e \cdot ||\bar{D}||$, where $\lambda_e$ is the electron wavelength ($\lambda_e = 0.025 \AA$ for 200 keV). Therefore, we can estimate that the tilt of the <0001> axis in the cone-shaped domain is $\theta \approx 0.6^\circ$. The slight tilt of the c-axis within the cone-shaped domain is most likely associated with strain relaxation. The misfit strain can be partially relieved through a tilt of the growth axis, similar to the process found for step terraces on miscut substrates.22,23,24 Given that the difference in lattice parameter c between AlN and GaN strained on AlN is approximately 0.275 Å, and the in-plane lattice parameter of AlN is 3.112 Å, relaxing the strain elastically in a unit cell would produce a maximum tilt of approximately 5°. Therefore, a tilt of 0.6° suggests that the local strain relaxation occurs over at least 3 nm in-plane. Since the QDs have a height of about 1.5 nm, the accumulated strain could relax over a large in-plane
distance, with the vertical correlation of the faceted structure propagating the tilt vertically. In fact, the 30° facets should give rise to nm-sized step terraces, depending on the actual step height, suggesting that relaxation could occur along these step terraces, similar to the mechanism described by S. Rathkanthiwar²². However, such a small tilt is not expected to contribute significantly to the redshift of the emission wavelength in these structures. Instead, alloy segregation and the different metal-to-nitrogen ratio and strain conditions in the cone-shaped domains lead to the formation of a secondary population of larger dots, responsible for the secondary redshifted emission.⁴

The APT analysis of a tip specimen containing cone-shaped domains is shown in figure 3. The animated version of this figure is supplied as Multimedia View. The reconstructed volume, using the Ga²⁺ and Ga⁺ ions presented in figure 3a, clearly shows the set of QDs, which appear as Ga-rich planes due to the limited resolution of the APT, with cone-shaped domains disrupting their planarity. The vertical resolution in La-APT for semiconductors is subnanometer and is typically better than the lateral resolution, though not as fine as in metallic samples. The lateral resolution for semiconductors in La-APT is generally poorer, typically around 1 nm at best, as determined for instance in SiC¹³. This limitation is caused by the perturbation of the ion trajectories at very close distances to the surface during field evaporation, which eventually affects the reconstruction of atom positions²⁵.

We will analyze in more detail four distinct regions throughout the reconstructed volume, which are identified by brackets in figure 3a: **R1**. A top region, where the QD planes are flat and homogeneous. **R2**. A second region containing a cone-shaped domain, primarily consisting of a core region with small, tilted boundaries. **R3**. A complex region, where the cone-shaped domain from **R2** is perturbed by a large, tilted boundary, likely belonging to a second cone-shaped domain only partially contained within the tip volume. **R4**. A region comprising the AlN buffer layer and its interface with the SL. A portion of a shallow pit can be observed at the interface, originating the second cone-shaped domain and its tilted boundary. Note that the Ga²⁺ signal present in the AlN layer displayed in figure 3a is noise from the tail of the N₂⁺ peak as presented in the supplementary material.

Figure 3b presents a CSR(Al) detector space map for each of the analyzed regions of the reconstructed volume. These maps are generated by integrating the counts over a depth of 10 nm along the z_s coordinate, corresponding to the blue boxes indicated in Figure 3a. The CSR(Al) map was chosen instead of the CSR(Ga) map because the barriers are composed of pure AlN, which results in unreliable CSR(Ga) maps in the barriers. As expected for III-N materials, a local increase in the CSR occurs at the <0001> pole in all of these maps.³,¹⁴,¹⁵ The pattern has nearly circular symmetry within the heterostructure region, with a CSR maximum at the <0001> pole growing from 12 in the CSR(Ga) map (i) up to 33 in the CSR(Ga) map (iv), corresponding to
maximum surface fields between 24.5 and 25 V/nm during the evaporation.⁸ The CSR(Al) maximum increases due to the modification of the tip shape during evaporation and, for the AlN buffer, due to the stronger bonding in the AlN layer compared to the SL, which contains weaker Ga-N bonds.¹⁴

<figure><img src="image_5.png" /><figcaption>FIG. 3. APT analysis. (a) Reconstructed volume of a tip specimen, displayed along the (x_s, z_s) - and along the (y_s, z_s)-plane (left and right panel, respectively). Only the Ga²⁺ (brown) and Ga⁺ (blue) ions are shown, to better visualize the presence of cone-shaped domains disrupting the planarity of the QD planes. (b) Detector space maps of the CSR(Al) extracted at the different reconstructed depths corresponding to the four blue boxes in (a). Note that the X and Y directions in detector space are parallel to the x_s and y_s directions in the reconstructed space, respectively. (c) Variation of the X (green) and Y (red) coordinates of the center of mass of the CSR(Al) in the detector space versus the reconstructed depth z_s. The animated version of this figure is supplied as Multimedia View.</figcaption></figure>

The evolution of the CSR(Al) center of mass coordinates versus the reconstructed depth z_s is presented in figure 3c. Moving along z_s from top to bottom, the CSR is initially slightly off centered at the beginning of the analysis but remains constant throughout R1. As z_s decreases and the topmost part of a cone-shaped domain is reached in R2, the center of mass gradually shifts towards the center of the detector and then remains constant until reaching R3. The total shift corresponds to about 0.84 cm in the detector space. The position of the CSR center of mass is primarily dependent on the orientation of the <0001> axis on the evaporating surface. These data suggest a modification in crystal orientation related to the presence of the cone-shaped domain, indicating that the core of the cone-shaped domain is misoriented. As z_s becomes more negative, the center of mass shifts again in R3, where the boundary of a second cone-shaped domain enters the evaporating volume. This shift appears to be clearly related to the boundary of the second cone-shaped domain, suggesting that both the core and boundary regions of the cone-shaped domains likely exhibit crystallographic misalignments. Note that the presence of differently oriented crystal
phases at the evaporating surfaces may locally deform the shape of the CSR(Al) detector space maps, as in the case of the CSR(Ga) map (iii), which is significantly broader and more asymmetric than the other patterns. It should also be noticed that other sources for this deformation, such as compositional variations, should be ruled out as all crossed regions (with the exception of the AlN substrate) are compositionally similar with Ga fractions in the QD planes in the range of 86-93%.⁴ Finally, the center of mass gradually drifts as the second cone-shaped domain occupies more of the evaporated volume, with the drift becoming more pronounced when the evaporation front reaches the AlN region in **R4**.

The angle $\Delta\theta(z_{s1}, z_{s2})$ between the orientations of the <0001> crystal axis at different specimen depths $z_{1s}$ and $z_{2s}$ can be approximated as:

$$ \Delta\theta(z_{s1}, z_{s2}) \approx \frac{|R_{<0001>} (z_{s2}) - R_{<0001>} (z_{s1})|}{L} $$

Where $L=74$ cm is the flight length of the atom probe. Considering the distance between the centers of mass in regions **R1** and **R2**, we determine the shift of the c-axis between a homogeneous region and a region that fully encompasses a cone-shaped domain with a tilted c-axis. This represents the misorientation of the c-axis within the cone-shaped domain. The norm of the vector connecting the centers of mass in regions **R1** and **R2** is $||\bar{D}|| = ||(0.6 \text{ cm}, 0.6 \text{ cm})|| = 0.84 \text{ cm}$. This corresponds to a maximum measured misorientation of the <0001> axis of $\Delta\theta(max) \approx 0.65^{\circ}$, which is consistent with the misorientation angle measured by CBED for a similar cone-shaped domain in another portion of the sample. These results demonstrate that the application of APT to the study of crystal orientation in III-N heterostructures is possible and can be carried out through a careful analysis of the distribution of charge states in the detector space. This methodology can be applied on samples containing extended defects impacting the orientation of nanoscale domains within the analyzed volume.

The most limiting factor regarding the resolution of this method for measuring the crystal orientation is that the nanoshape features must have lateral dimensions at least comparable to the tip diameter. Features smaller than the tip diameter result in an averaging of crystal-shifting effects, preventing conclusive interpretation. Additionally, the actual resolution limits of APT influence the crystallographic analysis. Vertical resolution is unlikely to significantly affect the interpretation of misorientation. Constructing detector space maps for checking crystal pole positions requires integrating over vertical dimensions larger than APT's typical resolution. However, limited lateral APT resolution could blur the detector space map, complicating the identification of pole axis orientations. Finally, it is important to emphasize that since the misalignment analysis occurs in the detector space, it is not affected by reconstruction artifacts. However, attributing the observed misalignment in detector space to nanoscale features requires accurate reconstruction. The method used in this study, though less spatially precise than electron
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diffraction and less accurate than X-ray diffraction for crystal orientation, offers two key advantages: it provides both crystallographic and chemical information simultaneously and enables the analysis of small volumes (50-100 nm in diameter) in 3D.

Surface reconstructions within the evaporated volume can also influence the determination of crystal orientation, as they can affect factors such as the local electric field or evaporation rates. However, this does not appear to be a significant factor for the sample studied in this work. Nonetheless, preferential evaporation could potentially impact the analysis near grain boundaries or in samples with significant compositional variation, as these factors might also influence surface reconstruction. While the method presented here has clear potential for application beyond nitrides, its robustness in complex semiconductor systems depends on several factors: the ability to detect the crystallographic axis across varying compositions, the structural similarity between phases in polycrystalline multiphase semiconductor systems, and the extent to which compositional gradients affect evaporation dynamics. Expanding the applicability of this technique to diverse material systems would likely require modifications to account for these variables. In such cases, Combining CBED and APT is promising, especially in light of recent advancements in APT integrated with TEM.²⁶ Correlating these techniques could overcome the limitations of CBED related to specimen projection and the resolution and surface-specific constraints of APT, providing valuable insights for interpreting crystallographic information in challenging semiconductor samples. This could pave the way for applying the method to even more challenging systems, such as polycrystalline alloys or complex heterostructures, where both structural and compositional challenges coexist.

In conclusion, we have explored the application of atom probe tomography to evaluate crystallographic misorientations within nanoscale semiconductor structures. The adventitious domains studied, appearing as cone-shaped formations emerging from shallow pits in AlN and extending through the AlGaN QD superlattices, reveal subtle but clearly resolved misorientation angles of their crystalline <0001> poles. This approach contributes to extending the possibilities of the already information-rich APT technique, underlining its capability to analyze crystallographic orientations within semiconductor nanostructures in addition to their composition and shapes.

In the supplementary material, the sample and cone-shaped domain structure and characteristic are presented in detail. The mass/charge spectra recorded in the gallium-rich layers containing the QDs and in the AlN buffer layer from the APT tip shown in figure 3 are presented. Also, the procedure used for the CSR(Al) center of mass determination in the detector space is detailed.
This project received funding from the French National Research Agency (ANR) via the ASCESE-3D and FUSSL projects (ANR-21-CE50-0016 and ANR-22-CE09-0024). 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).

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