Electric field-induced domain structures in ferroelectric AlScN thin films Open Access

The development and design of new non-volatile memory devices based on wurtzite-type ferroelectrics requires fundamental understanding of the individual switching mechanisms and their structure–property relationships linking the peculiarities of electrical data with detailed knowledge of the materials’ polar domain state.¹ Herein, the analytical capability of transmission electron microscopy (TEM) to investigate materials down to the atomic scale offers the potential to gain fundamental understanding of the electric bias dependent nanoscale morphology of polar domain structures, e.g., the characterization of the morphology of inversion domains and the present atomic structures of the domains and their domain walls.

Within the last couple of years, scanning (S)TEM investigations revealed the inversion of the unit-cell polarity in epitaxial $Al 1 − x Sc x N$ thin films on molybdenum bottom electrodes after ferroelectric switching² and were substantial to identify the reversal of the polar-direction after deposition of $Al 1 − x Sc x N$ films on GaN templates by sputtering.³ Moreover, our investigations showed the formation of inclined tail-to-tail inversion domain boundaries in single grains after a partial switching event. This demonstrated analog-type switching behavior in vertically strongly confined volumes.³ A breakthrough using in situ electron microscopy experiments was reported by Calderon et al., who succeeded to observe the mechanism of polarization switching ion of the atomic scale in $Al 1 − x B x N$⁠.⁴ Lately, the substantial progress in the deposition of high crystal quality thin films by metal–organic chemical vapor deposition (MOCVD) allowed one to study the electric field-induced domain pattern on larger scales.⁵ In the case of metal (M)-polar films, the preference for evolution of head-to-head domain walls after ferroelectric switching could be confirmed, thus implying that at least two fundamentally different types of domain walls can exist in wurtzite-type ferroelectrics.

The formation of charged domain boundaries within wurtzite-type ferroelectrics introduced previously not considered functionalities to this material class. The presence of mobile charge carriers to compensate the bound polarization charge in inclined head-to-head and tail-to-tail domain boundaries can allow the control of the films’ electrical resistance, leading to multilevel resistive switching depending on the applied voltage.^6,7 Herein, the unipolar states have higher resistance, whereas partial switching due to an applied voltage below the coercive field $E C$ enhances the density of inversion domain walls and, thus, leads to a variation of the carrier density inside the $Al 1 − x Sc x N$ films.

However, an observation-based description of the nanoscopic to atomic scale changes of the domain configurations and a generally accepted understanding of the interrelated atomic domain wall structures in $Al 1 − x Sc x N$ solid solutions and other wurtzite-type ferroelectrics yet remains elusive. So far, the necessary STEM studies could not be conducted due to limited film quality, which precludes the observation of domain patterns across larger areas. For instance, a columnar structure and a limited crystalline texture (⁠ $Al 1 − x Sc x N$ 0002 rocking curve full width at half maximum $>1 deg.$⁠) are in conflict with the dedicated STEM imaging conditions, which require high quality epitaxial thin films with low defect density. One promising strategy is the growth of lattice-matched films on GaN templates by precise control of the Sc content, which is non-consistently reported to be in the range of $0.09≤x≤0.14$ for non-ferroelectric samples grown by plasma assisted molecular beam epitaxy.^8–10 However, it is our general impression that structurally well-defined thin films typically show more complicated ferroelectric behavior than their fiber-textured and more defective sputtered counterparts possibly providing a reduced barrier for domain nucleation. Recently, we reported on advances in the growth of lattice-matched ferroelectric $Al 1 − x Sc x N$ (⁠ $x=0.13--0.15$⁠) thin films by MOCVD on GaN templates.⁵ The films showed epitaxial growth with superior crystalline coherence, which allows for an improved observation of structural details, such as nanoscale ferroelectric domain structures down to atomic resolution, and comprises a major advantage for structure investigations. In the previous work, we discussed the ferroelectric domain structure of these $Al 1 − x Sc x N$ films after switching from an all M-polar as-grown state to the electrically all nitrogen (N)-polar state after 400 switching cycles. The present study extends the description of atomic structures observed at field-induced domain walls and provides insights into stabilized domain structures within different addressed switched states.

A 230 nm thick ferroelectric $Al 0.85 Sc 0.15 N$ thin film sample was deposited onto a Si-doped n-GaN layer by MOCVD using a close-coupled showerhead reactor equipped with a proprietary setup for the generation of an adequate molar flow of the low-vapor pressure scandium precursor.¹¹ It was capped by a 10 nm thin layer of $SiN x$ to prevent oxidation. More details on the sample growth and electrode deposition can be found in our previous study.⁵ 100 nm thick Pt was deposited on top of the heterostructure via magnetron sputtering (Von Ardenne CS 730S), and top electrodes were structured via lithography and ion-beam etching. Electrical measurements were performed on circular electrodes with $20μ m$ in diameter using an AixACCT TF Analyzer 2000. All electrical measurements were performed using a voltage signal with a triangular waveform and a frequency of 1.5 kHz. The drive signal was applied to the top electrodes. Electron transparent cross-section lamellas were prepared by the focused ion-beam (FIB) method and extracted from the Pt/SiN/ $Al 0.85 Sc 0.15 N$/GaN capacitor structures orthogonal to the $⟨2 1 ¯ 1 ¯0⟩$ orientation.

Atomic scale investigation of the local atomic polarization was conducted on a probe $C S$-corrected JEOL NEOARM scanning transmission electron microscope operated at the acceleration voltage of 200 kV and a Thermo Fisher Spectra 300 instrument operating at 300 kV acceleration voltage. The annular bright-field (ABF)-STEM imaging mode was chosen for the identification of nitrogen atom positions using a collection angle-range of 10–20 mrad. Scan distortions and sample drift during image acquisition were minimized by fast serial recording of multi-frame images followed by post-processing image alignment. The rigid and non-rigid image registration of serial image stacks was performed using the Smart Align¹² (HREM Research, Inc.) plug-in running on the DigitalMicrograph v.3.5.1 (DM) (Gatan, Inc.) software. Fourier-filtering of non-rigidly processed ABF-STEM micrographs was applied using a simple radiance difference filter (lite version of DM plug-in HREM-Filters Pro/Lite v.4.2.1, HREM Research, Inc.) to remove high-frequency noise from the post-processed image. The local Sc composition was investigated by energy-dispersive X-ray spectroscopy (EDS) using a four-quadrant Super-X detector integrated into the column of the Spectra 300 instrument. These systems achieve highest count rates and are, therefore, capable of detecting small changes in chemical composition with highest sensitivity.^13,14 NBED recording was performed on the Spectra 300 microscope with an EMPAD detector. To enhance the accuracy of measuring both light and heavy atom positions, the differentiated differential phase contrast (dDPC) imaging method was used, which is based on evaluation of the center of mass (COM) signal recorded by a segmented Panther detector.¹⁵ 

As described in the previous work, the basal plane of n-GaN provides an epitaxial template for the $Al 0.85 Sc 0.15 N$ film, which is characterized by an out-of-plane c-axis (0002-)XRD rocking curve of 252 arcsec (0.07 deg.). Despite the high structural coherence, the epitaxial and single crystalline film diverges into a visible columnar structure in the upper half of the film having about 20 nm wide columnar grains, which are differentiated by the enhanced scattering contrast observed along vertical lines (see arrows in Fig. 1). These grain boundaries are consistent with the rough surface topology of the film. By EDS analysis, a minor Sc-enrichment of $x∼1$ at.% along the columnar grain boundaries is evidenced, which matches with similar observations on grain boundaries in biaxially oriented $Al 1 − x Sc x N$ films.¹³ Other features are apparent from their difference in scattering contrast, which show no measurable change of the local composition. These are attributed to slightly misoriented grains showing up at the film depth where the structural divergence starts and persistent inversion domains are observed, which are discussed in Sec. III B.

The study of the electric field-induced domain structure of the electrically all N-polar state was performed after the as-grown $Al 0.85 Sc 0.15 N$ film was subjected to 400 complete ferroelectric switching cycles and additionally to an unipolar voltage signal with a triangular waveform and an amplitude of 133 V to establish the investigated N-polar state. The corresponding current response is depicted in Fig. 2(b) (black). The larger number of switching cycles was necessary since the ratio of the required switching field $E C$ and the electrical breakdown field $E B$ as seen in polarization vs electric field (P-E) loops is evolving with a cycle number in these films, reaching stabilized conditions with up to 3500 cycles.⁵ Similarly, the current response for the positive branch changes to a typical ferroelectric displacement current peak after cycling for $≈$1000 times. More details on the ferroelectric properties of these films can be found in the previous work.⁵ 

Figures 2(a) and 2(c) present sketches of the polar domain structures and exemplary real space ABF-STEM images of the samples in its as-grown M-polar state and in the N-polar state after switching with $+133 V$ [Fig. 2(b)]. The formation of a strong zigzag shaped contrast has been described in the previous work showing an N-polar domain region at the top of the film and a pinned M-polar domain at the GaN interface with an inversion domain boundary (IDB) region in between both domains by atomic resolution imaging.⁵ The direct observation of field-induced domain patterns is enabled by the single-crystal like quality of the as-grown state, which essentially features fewer electron scattering centers resulting in homogeneous contrast across the complete layer. Indeed, the changes of scattering contrast in STEM images have been linked before to the presence of as-grown inversion domains, which is likely due to the perturbed electron channeling conditions along the viewing direction.^16,17 In this case, the three-dimensional geometry of domains with a suspected wedge-shape generates an overlap between multiple M- and N-polar domains along the finite lamella thickness, which explains the observed diffuse contrast pattern.

In this paragraph, we extend the investigations on the large-scale domain configuration by application of nanobeam electron diffraction (NBED) experiments in STEM (4D-STEM) and on the present atomic structures by atomic scale STEM analysis. In order to qualitatively connect the zigzag shaped scattering contrast in the presence of local perturbations of the undisturbed crystal lattice, e.g., the presence of polar domain boundaries and related electric fields, we applied the disk detection approach to NBED patterns to map the entire film area. The disk detection approach is based on the scan-point by scan-point acquisition and analysis of individual diffraction patterns obtained from nanometer sized areas.¹⁸ In principle, the analysis of an NBED-STEM pattern is able to reveal not only details on crystal orientation or strain, but also provides the possibility to retrieve information on local electric fields perpendicular to the incident electron beam direction.^19–21 For this purpose, the lateral shift of the central disk is evaluated by applying a cross-correlation algorithm.²² This method has not been yet applied to map out polar domains in wurtzite-type ferroelectrics.

The result of this analysis is depicted in Fig. 3, showing in (a) the ABF-signal from the investigated area and in (b) the corresponding electron beam shift in a color map. In agreement with the c-axis aligned spontaneous polarization direction in the wurtzite-type crystal structure, the calculated electron beam shift vectors feature a rotation by approximately $180 °$ (from purple to green color) at the location matching with the zigzag scattering contrast. The congruence of these signals indicates the presence of local structural features capable of deflecting the detected position of the central disk. One of the many reasons behind this could be the presence of local electric fields from domains with opposite polarity or other types of induced structural defects affecting the electron scattering conditions.

The atomic structure of electric field-induced inversion domain walls in wurtzite-type ferroelectrics is yet under debate. Therefore, the identified region showing inversion domains was investigated in detail by atomic imaging methods of aberration corrected STEM to investigate all present atomic structure details. Figure 3(c) displays an ABF-STEM image recorded within the boundary region. The yellow lines approximate the boundary of a spike domain after identifying the M-polar motif at its outer regions, whereas within the yellow marks, the projected atomic arrangement shows a dimer-like motif with pairs of vertically separated metal and nitrogen atomic columns. In a first approximation, a simple model with superimposed layers containing M- and N-polar unit cells could be constructed, which might explain the unrealistic short distances between the projected atomic column positions.

There exist several reports on the atomic structure of (impurity-stabilized) inversion domain walls in non-ferroelectric III-nitrides with respect to their horizontal, inclined, or vertical character. A selection of these observed structures is summarized in Fig. 4. As such, horizontal domain walls [Fig. 4(a)] are stabilized by anion-doping with oxygen forming extended and curved domain walls. Their structure is described by two interpenetrating N-polar and Al-polar wurtzite-type lattices, which share common anion sub-lattice positions forming a structural motif of triangular-atomic columns.^23,24 The identical motif [Fig. 4(b)] was observed for inclined domain walls with strong horizontal character by atomically resolved ADF-STEM.^17,25

In the case of the wurtzite-type structure, the vertical separation of domains with antiparallel $180 °$ orientation is associated with uncharged walls predicted by theory and for which the atomic positions were determined by TEM investigations of GaN and AlN.^25–29 The atomic structure of this domain wall is in principle described by a lateral unit-cell twin by which reverses the orientation of the hexagonal rings parallel to the (0001) plane and introduces large structural distortion by forming vertically stacked four- and eight-fold rings [cf. Fig. 4(c)]. Because of its calculated low-energy configuration, it is plausible that a similar low-energy atomic arrangement of the vertical domain wall is formed in wurtzite-type ferroelectrics as well.^30–33 

However, so far, the complicated domain structure in MOCVD grown wurtzite-type ferroelectrics impeded structural identification beyond the observation of supposed superposition structures at the inclined domain walls [Fig. 4(e)].⁵ We suppose that the projected structural motif consisting of triangular arranged atomic columns (termed “triangular atomic motif” in the following discussions) originates from the lateral overlap of two domains with antiparallel c-axis orientation in the projection direction inherent to a collective c-axis shift of the metal plane position in the inverted domain. The existence of such an overlap is a necessary consequence of the small horizontal domain size (10–20 nm diameter), which implies that the domain wall itself is curved with respect to the focal plane of the electron beam.

In delineation, a similar projected triangular atomic motif was discussed in connection to an intermediate state observed during polarization reversal experiments in ferroelectric AlBN.⁴ A possible structural model of this intermediate state was proposed based on the $β$-BeO structure [Fig. 4(d)], which is close to the parent wurtzite-type structure, but having an overall nonpolar supercell consisting of alternating Al-polar and N-polar dimers forming a projected structure with four- and eightfold rings, as in case of the vertical domain wall structure.

However, since a more or less diffuse triangular atomic motif was experimentally observed for all of the discussed domain wall models, it is very hard to determine whether the stabilization of an intermediate non-polar supercell [compare Fig. 4(c)] could have indeed a significant contribution to the overall contrast or if the M- and N-polar lattices on either side of a potentially atomically sharp domain wall are observed in superposition. Even if we would look perfectly parallel to a domain wall as it is the case in Figs. 4(a) and 4(b), we will still not be able to differentiate a domain wall with a supercell from a domain wall with two lattices in superposition. Therefore, only an edge-on view onto an atomically sharp vertical domain wall, if that is the electrically induced boundary structure, may refute speculations on a metastable supercell and provides new information on the real structure of electric field-induced domain walls in wurtzite-type ferroelectrics.

Next, we proceeded with the atomic structure investigations of the domain walls using the dDPC imaging mode. In Fig. 5, the growth of an N-polar spike domain into the persistent as-grown M-polar region which is extending 50–100 nm from the GaN interface is presented. Three regions providing magnified views onto the projected atomic structures and their simple interpretation by two-dimensional structure projections are given in panels (i)–(iii). The atomic arrangement depicted in panel (i) confirms the inverted N-polar atomic structure within the center parts of the spike domain, whereas panels (ii) and (iii) highlight regions at the edges of the spike domain. As shown in panel (ii), a clear edge-on view onto a vertical domain wall is allowed to image the atomic configuration at the domain interface formed after ferroelectric switching in $Al 0.85 Sc 0.15 N$⁠. The observed atomic motif is indeed consistent with the vertical low-energy structure between M and N-polar domains in GaN. Hence, it is confirmed that domain walls may appear atomically sharp when switching a wurtzite ferroelectric without the necessity to stabilize a potential nonpolar supercell. Moreover, potentially due to the curved nature of the domain wall, the projected superposition structure of the discussed triangular atomic motif is observed at the spike’s edges as well as the narrow region at the tip as highlighted in panel (iii).

The subtle differences in the projected dimer-like motif with vertically separated nitrogen and metal atom positions as observed within the highlighted region in Fig. 3(c) and the triangular atomic motif commonly found at the inclined domain walls [Fig. 5(iii)] are yet unexplained. The formation of atomically sharp interfaces, in this case a domain wall, is accompanied by local structural distortion of the parent lattice.^32,33 In the wurtizte-type nitrides, such distortions will impact the bond lengths and bond angles of M-tetrahedra. From the micrograph presented in Fig. 5, we determined the projected Al–N dumbbell angle in the clearly M-polar regions and clearly N-polar regions [panel (i)] and compared it to the angle at the inversion domain interface shown in panel (ii). Indeed, M-polar and N-polar regions exhibit an estimated dumbbell angle of $≈ 35 °$⁠, which is slightly reduced to $∼ 28 °$ at the two monolayers at the left side of the interface. Furthermore, a small vertical shift of metal positions is apparent at the interface. As a result, the projected atomic columns are displaced and, hence, will appear more diffuse in a three-dimensional superposition structure observed at the inclined domain walls.³¹ If this observed distortion is significant, the distortion of the M-tetrahedra will result in a local change of the internal lattice parameter u, which is defined as the vertical displacement between the metal and nitrogen atoms. Hence, the net polarization as well as the total energy barrier to invert the polarization is locally modified at the domain wall, potentially facilitating ferroelectric switching by domain wall movement. Recent calculations of local strain fluctuations in wurtzite-type ferroelectrics indicate huge impact on the switching pathways by allowing the system to reduce its switching energy barrier through a sequential switching pathway via metastable intermediate structures.^34–36 The capability to switch sequentially rather than taking the high energetic uniform switching pathway through the hexagonal-BN structure³⁷ suggests switching along individual lines of atoms, which can, in principle, allow extremely sharp domain walls.³² 

In summary, these observations provide detailed insights into the atomic structures observed at the boundary of the polarization inverted N-polar domains and the persistent M-polar interface domain. Despite formerly discussed projection effects of the wedge-shaped N-polar domains, our observations indicate that the electric field-induced domain walls can be indeed atomically sharp in $Al 0.85 Sc 0.15 N$⁠, as demonstrated by calculations.^32,33 This finding suggests that the formation of a metastable nonpolar supercell, e.g., of the discussed nonpolar $β$-BeO type supercell, is likely not generally necessary to stabilize a domain wall, albeit its existence cannot be ruled out completely due to the discussed projection issues and motivates further work on electric field-induced domain structures in wurtzite-type ferroelectrics.

After we have demonstrated a correlation of the enhanced zigzag shaped contrast in ABF images with the existence of wedge-shaped N-polar domains growing into the M-polar matrix, we can use the ABF contrast as a marker for the domain state of the material to investigate changes to the ferroelectric domain pattern after application of an external bias. In verification of this statement, as-grown capacitors were treated using the same voltage signal with a triangular waveform as before to cycle the material for 400 times into the above discussed N-polar state. From this state, we applied $−$153 V for switching to an electrically all M-polar state and $+$116 V (which is well below $E C$⁠) to achieve a partially switched domain state.

The resulting ferroelectric response and the corresponding domain states are summarized in Fig. 6, showing sketches and ABF-STEM images of the established domain patterns as well as the current response obtained when presetting the films into the respective state (compare Fig. S1 in the supplementary material).

For the switched electrically all M-polar state, the exemplary ABF-STEM micrograph [see Fig. 6(a)] displays a wedge-shaped domain pattern at the top electrode, but with a less pronounced contrast difference to the lower unipolar part of the film. Congruent to the location of the domains, atomic resolution investigations close to the top electrode interface [see Fig. 6(b)] reveal regions showing clear N-polarity, whereas regions below the N-polar domains equally show M-polarity.

This is the first time observation of stabilized domains with inverted polarity located at the top electrode after a backswitching event, whereas evidence for pinned layers at the substrate/bottom electrode interface has been reported before.^2,5 At this point, the question why it is easier to stabilize domains rather than to completely annihilate them remains speculation but is likely associated with either mobile carriers, which can screen the field or vacancies, which are easier to get into the film than out of the film. This result is consistent with earlier observations that there is a strong tendency for initial domain nucleation and switching of the initial minority polarization (in this case, N-polarity) to start at the upper electrode interface. Congruently, it switches back to the initial majority polarization from the pinned domain interface.^2,5,6 The presence of pinned domains has been postulated before to explain the initial imprint (i.e., the asymmetry of the P-E-loop around the polarization axis) and the occurrence of wake-up in certain wurtzite-type ferroelectrics.^30,38 Their direct observation in this work can, therefore, be considered to be an important contribution to understand the underlying mechanisms for both effects in this material class.

Upon a partial switching experiment [see Fig. 6(c)], the highly directed extension of the persistent inversion domains toward the GaN interface is observed. While the spikes seem to extend to a similar depth observed for the N-polar domains in the electrically M-polar state, after partial switching, the spikes appear thin and laterally intermittent distributed. Recent studies demonstrated the impact of local strain gradients within the film as originating from interfaces and chemical variations on the domain nucleation and switching pathways.^34,35,39 Hence, it seems very likely that ferroelectric switching is supported by the particular structural and chemical properties of this film with respect to the above discussed differentiation of the epitaxial constrained and highly coherent crystalline volume into a microstructure with local variations of Sc composition. As a particularity, no ferroelectric displacement currents are evident in the electrical response depicted in Fig. S1 of the supplementary material for reaching the partially switched state. Potentially, the horizontal extent of these spikes is small enough that no new screening charges have to be introduced and therefore, no ferroelectric displacement current can be detected on the positive branch. Further explanations on the electrical response are given in the supplementary material.

The direct observation of extended domains agrees with our common understanding on the progression of domain growth in ferroelectrics in general, where the domain wall motion along the thickness direction of a capacitor is typically much faster than the sideways motion.^30,40 Nonetheless, compared to the needle-like domains after partial switching, the sideways motion does obviously not occur homogeneously enough to result in columnar domains, but rather in wedge-shaped domains with charged domain walls. The formation of a zigzag polarization boundary and the wedge shape of the ferroelectric domains meeting the highly coherent lower crystal region could be energetically favorable over a uniform horizontal domain wall extending parallel to the (0001) wurtzite-type plane. Possible effects could be enhanced to charge screening properties of inclined domain walls, potential defects stabilizing the domain walls, or unfavorable bond-breaking at the $Al 1 − x Sc x N$/GaN interface. This observation is different compared to our previous STEM study on 4 nm thin $Al 1 − x Sc x N$ films, where we found evidence for purely horizontal domain walls. The modulation of the domain wall density and domain shape can be a potent strategy to enable new multi-bit memory devices.^5–7 We plan to elucidate these differences in more detail in our future work and motivate calculations of 3D structure models of the vertical domain walls in $Al 1 − x Sc x N$⁠.

The observation of electric field-induced domain patterns in Pt/ $Al 0.85 Sc 0.15 N$/GaN stacks is enabled by the correlation of atomic structure analysis and the modulated scattering contrast at the boundaries between those polar domains in ABF-STEM images. This allowed us to study the evolution of wedge-shape domains in dependence of the applied electric field. The ferroelectric domain patterns established after polarization reversal showed striking emergence of pinned domains. For the electric M-polar state, initial domains with N-polarity were observed at the top Pt electrode interface, whereas in the case of the electrical N-polar state, a persistent M-polar volume remains at the lower quarter of the film. The occurrence of pinning could be at least partially due to the covalent and directional bonding nature of the highly doped M-polar n-GaN interface, which results in excellent epitaxial crystallographic alignment of the initially grown defect-poor MOCVD $Al 0.85 Sc 0.15 N$ layer, but potentially increases resistance for complete polarization reversal. Atomic scale STEM analysis revealed that the atomic configuration of electric field-induced vertical domain walls in ferroelectric $Al 0.85 Sc 0.15 N$ can appear atomically sharp, consistent with structure calculations. This observation provides further experimental insight into the atomic configuration of polarization discontinuities in wurtzite-type ferroelectrics since their inclined and wedge-shape domain boundaries typically promote the observation of superposition structures in the projection direction. The designed engineering of polarization discontinuities, e.g., the inclination angle of the domain wall, would greatly affect its properties, such as the density of bound charge, and could potentially lead to the development of novel devices for in-memory computing.

See Fig. S1 in the supplementary material that adds additional information on the discussion of electrical data with backswitching from the partially switched state.

This collaborative work was enabled through funding by the Federal Ministry of Education and Research (BMBF) under Project No. 03VP10842 (VIP $+$ FeelScreen) and in project ProMat_KMU “PuSH” Grant No. 03XP0387B and the Deutsche Forschungsgemeinschaft (DFG, German Research Foundation)—Project ID 434434223—SFB 1461; Project ID 286471992—SFB 1261 as well as Project ID 458372836 and Project ID 448667535. The TEM in the group of AR was funded by the Deutsche Forschungsgemeinschaft (DFG, German Research Foundation) –DFG Project No. INST 144/462-1 FUGG. This work was also funded by the European Union (FIXIT, GA 101135398). Views and opinions expressed are, however, those of the author(s) only and do not necessarily reflect those of the European Union or the European Research Council Executive Agency. N.W. thanks Christin Szillus for TEM sample preparation using the FIB technique.

The authors have no conflicts to disclose.

Niklas Wolff: Conceptualization (lead); Investigation (equal); Visualization (equal); Writing – original draft (lead). Tim Grieb: Formal analysis (lead); Investigation (equal); Software (equal); Writing – review & editing (equal). Georg Schönweger: Investigation (supporting); Writing – original draft (supporting). Florian F. Krause: Formal analysis (equal); Software (equal); Writing – review & editing (supporting). Isabel Streicher: Resources (equal). Stefano Leone: Funding acquisition (equal); Resources (equal). Andreas Rosenauer: Resources (equal); Writing – review & editing (supporting). Simon Fichtner: Funding acquisition (equal); Writing – review & editing (equal). Lorenz Kienle: Funding acquisition (equal); Resources (equal); Writing – review & editing (equal).

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