# Interplay between Crystal Structure and Optical Response in Plateau–Rayleigh Zn<sub>2</sub>GeO<sub>4</sub>/SnO<sub>2</sub> Heterostructures

Jaime Dolado, Fernanda Malato, Jaime Segura-Ruiz, Ruth Martínez-Casado, María Taeño, Irina Snigireva, Pedro Hidalgo, Bianchi Méndez, Gema Martínez-Criado

► To cite this version:

Jaime Dolado, Fernanda Malato, Jaime Segura-Ruiz, Ruth Martínez-Casado, María Taeño, et al. Interplay between Crystal Structure and Optical Response in Plateau–Rayleigh Zn<sub>2</sub>GeO<sub>4</sub>/SnO<sub>2</sub> Heterostructures. Advanced Photonics Research, 2023, 4 (8), pp.2300063. 10.1002/adpr.202300063. hal-05079993

HAL Id: hal-05079993

https://hal.science/hal-05079993v1

Submitted on 22 May 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
# Interplay between Crystal Structure and Optical Response in Plateau–Rayleigh Zn<sub>2</sub>GeO<sub>4</sub>/SnO<sub>2</sub> Heterostructures

Jaime Dolado,\* Fernanda Malato, Jaime Segura-Ruiz, Ruth Martínez-Casado, María Taeño, Irina Snigireva, Pedro Hidalgo, Bianchi Méndez, and Gema Martínez-Criado

Hybrid semiconductor nanowire (NW) heterostructures are an ideal playground for cutting-edge optoelectronic nanodevices. Among the several synthesis methods, Plateau–Rayleigh (PR) crystal growth is an effective route for producing decorated NWs with unprecedented properties. However, lateral variations in composition and/or crystal structure are postulated to play a central role in their optical response, but it is difficult to probe correlatively the elemental order, atomic organization, and light emission on length scales of tens to hundreds of nanometers. Usually, electron microscopies are applied to address the formation of clusters and imperfections in representative cross sections of the samples. Herein, a simultaneous spatially resolved nano-analysis of the crystal symmetry, chemical composition, and optical properties of a whole Zn<sub>2</sub>GeO<sub>4</sub>/SnO<sub>2</sub> NW heterostructure produced by PR instability is provided. The observations show the connection among Zn impurities, secondary phases, and asymmetrically distributed UV emissions present in the crystallites decorating the NW (Zn-doped Sn<sub>1-x</sub>Ge<sub>x</sub>O<sub>2</sub>). The contributions of the elemental diffusion, crystal domains, and atomic site configurations to the light-emission phenomena are disentangled in these hybrid architectures. The findings elucidate unknown underlying mechanisms that are critical to tailor emergent properties for rationally designing novel complex nanodevices based on 1D materials.

## 1. Introduction

Decorated semiconductor nanowires (NWs)[1] are promising architectures for next-generation nanoelectronic devices. Their widely tunable electronic,[2] optical,[3] and plasmonic[4] capabilities make them desirable hosts for a wide range of quantum,[5] sensing,[6] and energy-harvesting[7] applications. By manipulating the growth kinetics, it is now possible to produce extremely complex nanoscale systems. New electrical and photonic functionalities are generated by the addition of sequential branches with differences in composition and/or doping.[8–10] Among the several synthesis methods, the Plateau–Rayleigh (PR) crystal growth enables the formation of decorated NW systems. Unconventional morphological characteristics with modulations in periodicity, size, as well as cross-sectional anisotropy are produced by controlling the deposition conditions.[11,12]

Within the semiconductor families, wide bandgap (WBG) oxides, such as SnO<sub>2</sub> (E<sub>g</sub> ≈ 3.6 eV) and Zn<sub>2</sub>GeO<sub>4</sub> (E<sub>g</sub> ≈ 4.5 eV), exhibit relevant properties: visible-light transparency, high conductivity (≈10<sup>4</sup> S cm<sup>-1</sup>), and large electron mobility (≈80 cm<sup>2</sup> V<sup>-1</sup> s<sup>-1</sup>).[13] High-power electronics, ultraviolet (UV) photodetectors as well as UV light-emitting diodes are just a few applications that can profit from the substantial breakdown electric field and short-wavelength emission induced by the wide energy bandgap.[14] Also, interfacial charge transfer for CO<sub>2</sub> reduction is possible by band structure engineering of WBG oxides.[15]

Therefore, the application of PR growth approach to WBG oxides represents an excellent pathway to further expand the potential of hybrid NW heterostructures for advanced semiconductor nanotechnologies. Here, we report a correlative study of a single Zn<sub>2</sub>GeO<sub>4</sub>/SnO<sub>2</sub> NW architecture formed by PR-instability mechanism using a synchrotron nanoprobe. Since the performance of the heterostructure is intimately related to the crystal quality, it is crucial to examine their physical properties at the nanoscale. Commonly, native defects, heterogeneities, and/or structural imperfections drastically affect the optoelectronic responses.[16–18] So far, energy dispersive X-ray spectroscopy (EDS) and correlative high-resolution transmission electron microscopy (HRTEM) have been mostly used to assess the

<footnote>J. Dolado, J. Segura-Ruiz, I. Snigireva, G. Martínez-Criado
Experiments Divison
X-ray Nanoprobe Group
European Synchrotron Radiation Facility
38043 Grenoble, France
E-mail: doladofe@esrf.fr</footnote>

<footnote>F. Malato, G. Martínez-Criado
Instituto de Ciencia de Materiales de Madrid
Consejo Superior de Investigaciones Científicas
28049 Cantoblanco, Spain</footnote>

<footnote>R. Martínez-Casado, M. Taeño, P. Hidalgo, B. Méndez
Departamento de Física de Materiales
Facultad de Ciencias Físicas
Universidad Complutense de Madrid
28040 Madrid, Spain</footnote>

<footnote>The ORCID identification number(s) for the author(s) of this article can be found under https://doi.org/10.1002/adpr.202300063.</footnote>

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

<footnote>DOI: 10.1002/adpr.202300063</footnote>
structure and composition of hybrid semiconductor NW systems.[19–23] Local optical heterogeneities have also been identified using indirect techniques like cathodoluminescence.[24,25] However, a thorough and in-depth chemical and structural examination of the overall system is crucial. Although the creation of crystalline phases driven by PR mechanism has been proved, there are still many open questions. For example, the role of impurities on coupling formation and local atomic site configuration,[26] how these necklace arrangements are joined to form decorated architectures,[27–29] whether there is a full control over composition, diffusion paths, and/or structural modifications,[10,11,30,31] coexisting crystalline polytypes,[20,32] diffusion-driven effects,[33] phase separation,[34] radial or axial chemical modulations,[20,35,36] as well as atomic scale variability due to fluctuations in the growth rate.[10,37]

In the present work, we address some of these issues in a representative Zn₂GeO₄/SnO₂ NW architecture produced by PR method.[38] For the first time, the elemental composition, structural order, and optical response of these decorated NW systems are simultaneously investigated at the nanoscale using X-ray fluorescence (XRF), X-ray-excited optical luminescence (XEOL), X-ray diffraction (XRD), and X-ray absorption spectroscopy (XAS). This “all-in-one” strategy, which makes use of a hard X-ray nanobeam, provides a complete correlative image of a single-decorated NW. For our experiment, we used the European Synchrotron Radiation Facility (ESRF)’s nano-analysis beamline ID16B.[39]

**2. Results and Discussion**

**2.1. Radial and Axial Elemental Modulation**

The analysis of the elemental composition in decorated NWs is a
crucial step in the characterization of these hybrid systems, as it
provides information on the type and number of dopants, as well
as the heterogeneities present in the material. **Figure 1a** shows
the scanning electron microscopy (SEM) image of a single PR
heterostructure. It consists of a straight central wire (5 µm length
and 100 nm diameter) surrounded by well-faceted crystallites
(sizes between 500 and 800 nm). Figure 1b shows the XRF nano-
imaging results obtained by raster-scanning the Zn₂GeO₄/SnO₂
heterostructure in the X-ray nanobeam, where the superimposed

distributions of Zn (blue), Ge (green), and Sn (red) are displayed
(average XRF spectrum and individual elemental distribution
maps can be found in Figures S1 and S2, Supporting
Information, respectively). Although it is predicted that the
NW axis mostly consists of Zn₂GeO₄, the XRF maps of the crys-
tallites show a clear and interesting sandwich-like arrangement.
The creation of an alloy that may significantly alter the structural
order is suggested by the considerable inclusion of Ge in the
SnO₂ crystallites. However, there is no evidence of a one-to-
one connection between the Ge and Sn distributions. The Ge
atoms are preferentially located within the conical faceted sides
of the crystallites (sandwich-like structure compared to Sn),
which may result in the creation of secondary phases or defects
and consequently impact the spatial distribution of the carrier
recombination processes. Moreover, Sn localization inside crys-
tallites is proportional to the crystal volume, with a higher inten-
sity in thick, fully faceted portions, whereas Ge is distributed in
the regions that are not totally faceted (conical sides). Finally, in a
minor fraction, the Zn signal comes mainly from the NW and
slightly from the conical-faceted sides of the crystallites (see
Figure S2, Supporting Information). In brief, our XRF findings
spot the massive incorporation of Ge into the SnO₂ crystallites
(accompanied also by Zn in a lower degree), suggesting
Sn₁₋ₓGeₓO₂ alloy formation with a particular sandwich-like
spatial pattern.

In good agreement with these findings, we recently reported
the observation of a high Ge quantity within the SnO₂ crystal-
lites (from 8% to 17%) by TEM-EDS.[38] Our results point to
the formation of a Sn₁₋ₓGeₓO₂ solid solution, which preserves
the rutile crystal structure. Likely, the faceted areas present a
lower amount of Ge (and a higher amount of Sn) because
the Ge out-diffusion to the Zn₂GeO₄ NW surface (which is
arranged in periodic droplets due to PR-instability) acts as
nucleation sites for the Sn crystallites. The more faceted the
Sn₁₋ₓGeₓO₂ crystallites is, the more Sn replaced Ge to form
the rutile crystal phase. Recent studies have demonstrated that
the NW coating in the PR-type processes is dictated by the evo-
lution of the internal crystalline nanostructure over time, or by
the exchange of free atoms between the NW surface and the
enveloping near-surface layer.[20,32] Therefore, these factors
could be involved in the different distribution of Ge and Sn
within the PR heterostructure, determining its physical
properties.

<figure><img src="image_2.png" /><figcaption>**Figure 1.** Elemental distribution obtained by nano X-ray fluorescence (nano-XRF). a) Scanning electron microscope (SEM) image of Plateau-Rayleigh (PR) heterostructure. b) XRF map in RGB (Red, Green, Blue) visualization that depicts the XRF intensities of Sn (red), Ge (green), and Zn (blue). Their color brightness (light represents high counts, dark low counts) indicates the intensity ranges (more details can be found in Supporting Information).</figcaption></figure>
**2.2. Anisotropic Optical Response**

The optical response of PR heterostructures is a critical parame-
ter that determines their functionality and suitability for optoelec-
tronic applications. Here, we have made use of the optical
contrast created by XEOL to concurrently capture spectrally
and spatially resolved luminescence signals and XRF data for
each sample point (see **Figure 2**). Following primary XAS, the
majority of photogenerated electrons and holes recombine
radiatively via secondary XRF and XEOL pathways. Figure 2a
presents the average XEOL spectrum, where a main emission
dominates in the visible range with two components located
at 2.0 and 2.4 eV. It should be noted that the prominent peak
exhibits multiple resonances overlapping the main bands, which
can be attributed to light reflections between the parallel facets of
the $Sn_{1-x}Ge_xO_2$ crystallites.[³⁸] In addition, there is a weak band
in the UV range located at 3.3 eV. The dominant band corre-
sponds to the emission related to the transition of the trapped
electrons from oxygen vacancies to intrinsic surface states from
the $Sn_{1-x}Ge_xO_2$ crystallites. The luminescence of $SnO_2$ is
extremely complex, exhibiting a broad-emission band in the visi-
ble range, consisting of three components (orange, green, and
blue at 2.1, 2.5, and 2.7 eV, respectively) associated to oxygen
vacancies.[⁴⁰,⁴¹] These components were also reported from
$SnO_2$ nanoribbons by XEOL excited with both soft and hard
X-rays.[⁴²,⁴³] Regarding the UV band, its origin is uncertain.
This emission could come from the $Zn_2GeO_4$ NW, which exhib-
its a wide emission associated to oxygen vacancies in the UV
range.[⁴⁴] However, an UV emission related to surface states
exhibited by $SnO_2$ nanorods and nanoparticles has also been
reported,[⁴⁵] so it could also come from $Sn_{1-x}Ge_xO_2$ crystallites.

Figure 2b,c displays the spatial distribution of the respective XEOL bands on the nanometer scale. The visible emission associated to oxygen vacancies is mostly distributed within the $Sn_{1-x}Ge_xO_2$ crystallites, which becomes especially stronger where Sn signal is higher. In contrast, the less intense UV band of unidentified origin is only emitted from one side of the $Sn_{1-x}Ge_xO_2$ crystallites (for more clarity, see Figure S4, Supporting Information), which raises the possibility of an orientation-dependent phenomenon connected to the crystal structure. Earlier investigations have reported faceted dependent cathodoluminescence in $SnO_2$ microtubes, which would imply different distribution of impurities and/or defects.[⁴⁰,⁴⁶] Alternative explanations for the UV emission include

preferential incorporation of defects coming from the out-
diffusion of Ge or Zn from the $Zn_2GeO_4$ NW, which would
produce electronic states in the energy bandgap and cause the
UV emission. The inhomogeneous variation of the radiative cen-
ters could probably be associated with the different distribution
of impurities and/or structural imperfections. The luminescence
bands in semiconductors are widely known to be tightly corre-
lated with the density of defect-related energy levels.[⁴⁵,⁴⁷]

**2.3. Correlation between Optical Response and Elemental Distribution**

It is possible to determine how the spatial distribution of heter-
ogeneities affects the optical response by combining spatially
resolved XRF and XEOL methods. **Figure 3** represents the nor-
malized radial line profiles obtained along the PR heterostructure
(see Figure 3a). Axial line profiles are also shown in Supporting
Information (Figure S5, Supporting Information). Figure 3b–d
shows the XRF and XEOL profiles data obtained from each of
the three $Sn_{1-x}Ge_xO_2$ crystallites along the perpendicular direc-
tion to the core $Zn_2GeO_4$ NW. The Sn signal (red lines) has a
maximum intensity near the center of the $Sn_{1-x}Ge_xO_2$ crystalli-
tes and follows a similar trend in all cases. On the contrary, the
Ge signal (green line) of the largest crystallite presents a different
profile. In general, a tiny portion of the Ge-related photons could
suffer self-absorption effects, which may marginally increase Zn
and Sn signals. An increased Ge intensity can be seen coming
from the structure’s edges where the biggest crystallite looks to
be entirely created. Contrarily, the smaller $Sn_{1-x}Ge_xO_2$ crystal-
lites exhibit a more uniform Ge incorporation, with a profile
resembling a Gaussian in accordance with their morphological
shape. The Zn profiles (blue lines) also show a multipeak shape,
which denotes a rather unusual and distinct behavior inside the
crystallite volume. The Zn signals from the crystallite and the
NW overlap in the lines can center because the $Zn_2GeO_4$ NW
is the source of the majority of the contribution. As a result,
the findings not only demonstrate Zn’s integration into
$Sn_{1-x}Ge_xO_2$ crystallites but also its preferred spatial location
toward the crystallite’s edges. The dependency of the crystallo-
graphic orientation on the incorporation of dopants and impuri-
ties, which causes different elemental diffusion and trapping
with subsequent material deposition, may be the cause of this
Zn distribution.[⁴⁸,⁴⁹]

<figure><img src="image_3.png" /><figcaption>**Figure 2.** X-ray-excited optical luminescence (XEOL) distribution with nanometric resolution. a) Average XEOL spectrum recorded over the PR hetero-structure shown in Figure 1a. b,c) XEOL-integrated intensity maps over the energy ranges defined by the colored bands (visible and UV) in (a).</figcaption></figure>
<figure><img src="image_4.png" /><figcaption>**Figure 3.** Direct link between luminescence emissions and elemental distribution. a) XRF nanoimage of PR heterostructure on which the line profiles were collected. b–d) Normalized XRF line profiles (top) for Sn (red symbols), Ge (green symbols), and Zn (blue symbols), andXEOL line profiles (bottom) for visible (orange symbols) and UV (purple symbols) emissions. All profiles were collected along the white-dashed lines drawn on the PR heterostructure shown on the left-hand side, perpendicular to the central nanowire (NW).</figcaption></figure>

In addition to probing different sample depths (XEOL is considerably more surface sensitive than XRF), the radial line profiles of the XEOL signals add additional details about the three $Sn_{1-x}Ge_xO_2$ crystallites. In contrast to the XRF signal, the spatial distribution of XEOL is not only governed by the beam spot size and penetration of the X-ray nanobeam within the sample (energy dependent, leading to a finite-generation volume), but also by the diffusion of the photogenerated carriers. Several elements might have an influence on the resulting light profile like sample morphology, composition, internal absorption, and/or reflection mechanisms, as well as secondary photon emissions. The visible band profile (orange line), which has a Gaussian-like shape and is strongly associated with Sn, follows the crystal volume (albeit it has a reduced full width at half maximum owing to diffusion effects[50]). The UV-emission profiles show a much thinner line width peaking near the left edge of the $Sn_{1-x}Ge_xO_2$ crystallite, which matches one of the maxima in the Zn profile nicely. Axial line profiles taken along the PR heterostructure (Figure S5, Supporting Information) further illustrate the direct relationship between the UV band and Zn localization. These XEOL results could indicate a significant emission polarization anisotropy of the UV luminescence due to likely the faceted nature of the $Sn_{1-x}Ge_xO_2$ crystallites linked to the $Zn_2GeO_4$ NW.

In short, the correlative XRF-XEOL analysis shows further Ge and Zn element colocalization in the sandwich-like regions of $Sn_{1-x}Ge_xO_2$ crystallites. In addition, for all the cases, the UV emission is closely related to the Zn distribution, which is mostly concentrated at one of the edges of the $Sn_{1-x}Ge_xO_2$ crystallites. This observation suggests that the anisotropic UV emission could come from recombination centers coupled to complexes associated to the presence of Zn.

## 2.4. Secondary Phase Formation

In addition to chemical composition, secondary phases that result from changes in crystalline structure have a significant

impact on the optical characteristics and optoelectronic performance of materials. Thus, to investigate the long-range structural order of the PR heterostructure, nano-XRDs were also performed on the $Sn_{1-x}Ge_xO_2$ crystallites using the hard X-ray nanobeam. TEM measurements were previously used to pinpoint the rutile structure of the $Sn_{1-x}Ge_xO_2$ crystallites. When seen along [110], the facets of the $Sn_{1-x}Ge_xO_2$ crystallites are truncated by (001) planes (see Figures S6 and S7, Supporting Information). **Figure 4b** shows three representative XRD patterns of the (110) reflection, that were collected over three different zones of the PR heterostructure (see Figure 4a). A close inspection of the reflection (110) in the center (point 2) and one of the edges of the $Sn_{1-x}Ge_xO_2$ crystallite fully formed (point 3) reveals a Gaussian-like XRD peak at about 3.32 Å for both areas with similar linewidth, indicating the presence of a uniform crystalline alloy in these regions. However, there are two Gaussian contributions at around 3.31 and 3.33 Å on the opposite side of the crystallite (point 1), indicating the presence of two distinct $Sn_{1-x}Ge_xO_2$ phases. According to the first-principles calculations shown in Table S1, Supporting Information, the shift of the XRD peak to 3.33 Å may be connected to a lower concentration of Ge in the alloy.

Figure 4c,d displays the spatial distribution of the respective XRD phases on the $Sn_{1-x}Ge_xO_2$ crystallite. The $Sn_{1-x}Ge_xO_2$ crystallite's volume is essentially covered by the phase that peaked at ≈3.31 Å. However, the $Sn_{1-x}Ge_xO_2$ crystallite's phase peaked at 3.33 is predominantly present on one side and exhibits a spatial distribution in the crystallite that is extremely comparable to that of the UV emission (see XEOL map, Figure 2c). Given that the regions of the $Sn_{1-x}Ge_xO_2$ crystallite with the lowest 3.31 Å phase signal and the greatest 3.33 Å phase signal coincide, the signals of the two phases are complementary. The UV emission is therefore directly observed in the region of the PR heterostructure where the XRD results show a change in the long-range structure. The formation of the secondary phase
<figure><img src="image_5.png" /><figcaption>**Figure 4.** Revealing a secondary phase at the nanoscale. a) XRF nanoimage of the PR heterostructure. b) X-ray diffraction (XRD) pattern of the (110) reflection (solid circles) acquired at three different areas of the sample, indicated by numbers in (a). Different Gaussian contributions are plotted with dotted lines. c,d) XRD-integrated intensity maps over the *d*-spacing ranges defined by the colored bands (≈3.31 and ≈3.33 Å) in (b).</figcaption></figure>

within the Sn₁⁻ₓGeₓO₂ crystallites might possibly be explained by a Zn impurity-assisted mechanism, which would have an impact on the degree of optical light and strain generated.

## 2.5. Zn Local Environment in the PR Heterostructure

Given the key role of Zn atoms within the Sn₁⁻ₓGeₓO₂ crystallites, nano-X-ray absorption near-edge spectroscopy (nano-XANES) data were also taken around the Zn K-edge to get a deeper insight into its local structure. **Figure 5a** shows the spatially resolved XANES acquired over the three same areas where the nano-XRD measurements were collected. The XANES signal is analogous to the partial density of the empty states of the absorbing atoms. In particular, XANES collected at the Zn K-edge reveals the *p*-partial density of states in the conduction band. Here, two spectral XANES regions can be distinguished: a prominent white line assigned mainly to the 1s → 4p dipolar transitions, and the post-edge resonances, which are due to multiple scattering effects. All nano-XANES spectra show an intense

narrow peak at 9670.2 eV attributed to the white line. The post-edge oscillations are quite similar in the sample areas marked by blue and green dots, indicating a similar local environment around Zn atoms. However, the smoother wiggles present in the XANES collected at the sample area indicated by the point 1 could be associated to the extra structural disorder caused by the presence of two crystal phases detected by XRD (Figure S8, Supporting Information).

One popular and empirical assumption is that the XANES spectrum from an unidentified local structural order may be interpreted as a linear superposition of the XANES spectra from two or more well-known samples. As a result, Figure 5b plots the XANES spectra of three model compounds, in addition to one acquired from the Sn₁⁻ₓGeₓO₂ crystallites. In ZnO, Zn atoms are in tetrahedral sites.[51] In the Sn₀.96Zn₀.04O₂ alloy, Zn is introduced into the SnO₂ structure with no major change in the crystal structure.[52] Finally, the XANES spectrum around Sn K-edge acquired in SnO₂ rutile type is shown as well,[53] where each tin atom is expected to be coordinated by six oxygen atoms (octahedral coordination). Our results strongly suggest a substitutional

<figure><img src="image_6.png" /><figcaption>**Figure 5.** Zn local environment in the PR heterostructure. a) X-ray absorption near-edge spectroscopy (XANES) data around the Zn K-edge taken in the same three sample areas where XRD measurements were acquired (see Figure 4a). b) XANES data recorded around the Zn or Sn K-edges of reference materials with tetrahedral (ZnO[51]), distorted octahedral (Sn₀.96Zn₀.04O₂[52]), and octahedral (SnO₂[53]) coordination, in addition to the one acquired in the PR heterostructure (at the bottom of the graph). The spectra were shifted vertically for clarity.</figcaption></figure>
<figure><img src="image_7.png" /><figcaption>**Figure 6.** a) Sketch illustrating the PR-growth of Zn₂GeO₄/Sn₁₋ₓGeₓO₂ decorated NW architecture. b) Suggested energy diagram for theXEOL process in Sn₁₋ₓGeₓO₂ and Zn-doped Sn₁₋ₓGeₓO₂. Blue and white circles represent electrons and holes, respectively. Defect states are represented by horizontal lines in the bandgap.</figcaption></figure>

Zn incorporation into the Sn sites. A direct comparison of XANES data from the post-edge region around the Zn and Sn K-edges for the PR heterostructure and SnO₂, respectively, indicates the predominant octahedral coordination of Zn atoms in the Sn₁₋ₓGeₓO₂ structure. Our findings are corroborated by theoretical ZnO in a rutile-like environment XANES simulations that display comparable spectrum patterns.[54] Thus, the fact that the XANES spectra of the Sn₁₋ₓGeₓO₂ crystallites is significantly different from the XANES data of the Sn₀.96Zn₀.04O₂ alloy also supports that Zn is incorporated as a dopant into the rutile structure (Sn sites). In addition, the white line of the PR heterostructure is shifted to higher energies with respect to those of ZnO and the Sn₀.96Zn₀.04O₂ alloy, indicating an increase in Zn coordination.

In short, XANES data indicate a substitutional incorporation of Zn dopants into Sn sites in the rutile crystal lattice with octahedral coordination inside the Sn₁₋ₓGeₓO₂ crystallites. These Zn-related defects might promote phase-separation phenomena, as well as electronic states within the bandgap responsible for the UV emission observed by XEOL.

## 2.6. Full Correlative Picture of the PR Heterostructure

According to the elemental nano-analysis previously described, the following sequence for the PR growth of the decorated NWs can be drawn (Figure 6a): 1) thermal growth of the precursors leads to the formation of Zn₂GeO₄ NWs, on which an array of amorphous GeO₂ (a-GeO₂) droplets along the NW is initially produced via the PR instability; 2) the a-GeO₂ droplets act as preferential sites for Sn nucleation, giving rise to Sn₁₋ₓGeₓO₂ crystallites that encapsulate the NW; 3) the more faceted the Sn₁₋ₓGeₓO₂ crystallite is, the more Sn replaced Ge, which leads to the Ge/Sn sandwich pattern.

In addition, a 3D rough picture of the interplay between crystal structure and optical response in the decorated NW can be obtained through the “all-in-one” strategy. Figure 6b shows a schematic representation of the potential energy band diagram resulting from all our results. The Zn-doped Sn₁₋ₓGeₓO₂

crystallite exhibited an asymmetrically distributed UV emission, which has been collected only in the areas with the highest Zn impurity content, suggesting the existence of new electronic states assisted by Zn complexes. The XRD results support the presence of a secondary crystalline phase right in the areas with the highest amount of Zn, which enter into the crystal lattice replacing Sn as indicated by XANES results.

## 3. Conclusion

Due to the complexity of the multiphase chemical phenomena involved at the nanoscale, it is difficult to understand the performance of novel decorated NW heterostructures produced via PR instability. In this work, we have characterized a representative Zn₂GeO₄/SnO₂ NW architecture, which has been produced by the PR instability. By combining multiple synchrotron nano-analysis techniques (nano-XRF, nano-XEOL, nano-XRD, and nano-XAS), we shed light on the origin of the UV emission distributed asymmetrically across the heterostructure. Our findings reveal a connection between the Zn species and the UV radiative transitions. In addition, nano-XRD data show the formation of an additional phase, which is well correlated with the UV emission. To close, nano-XANES measurements indicate an octahedral coordination of Zn atoms in the Sn₁₋ₓGeₓO₂ crystallites. Thus, our results suggest that a Zn impurity-assisted process could be responsible for the UV optical emission, involving also the formation of a secondary phase within the Sn₁₋ₓGeₓO₂ crystallites. Our work showed that a multimodal synchrotron nanoprobe is an effective “all-in-one” method for obtaining a thorough understanding of intricate nanosized hybrid systems at the atomic and electronic levels.

## 4. Experimental Section

**Synthesis:** The syntheses of the NW heterostructures were carried out by a single-step thermal evaporation method. The material precursors were a compacted mixture of ZnO:Ge:C powders (2:1:2 wt%), where SnO₂ powders were added in a proportion of 10 wt% related to the total
ZnO:Ge amount. These precursors acted as source and substrate, with no foreign catalyst. The thermal treatments were conducted at 800 °C for 8 h under an Ar flow of 1.5 L min⁻¹ (see more details in Supporting Information).

**X-ray Characterization:** The hard X-ray nanoprobe of the nanoanalysis beamline ID16B of the ESRF was used in our study. Schematic of the experimental setup for the multiple X-ray techniques at the nanoscale is shown in Figure S10, Supporting Information. The characteristics XRF is recorded with an energy dispersive Si drift detector at an angle of θ = (15 ± 5)° with respect to the sample surface. The XEOL is detected normal to the sample surface by a far-field optical collection system and in backscattering configuration, which avoids lateral shadowing effects (see Figure S11, Supporting Information). The monochromatic X-ray beam was focused to 80 × 60 nm² spot size (see Figure S12, Supporting Information) with about 10¹¹ ph s⁻¹. The XRF and XEOL maps were taken at 11.25 keV (excitation energy above the Ge K-edge located at 11.1031 keV) with a pixel size of 50 × 50 nm² over a 3 × 5 µm² sample area and 1 s/point integration time. For XRD measurements, the Sn₁₋₄Ge₅O₂ crystallites were illuminated by a 14 keV monochromatic nanobeam and scanned with a pixel size of 50 × 50 nm² and 0.5 s per point integration time. The XRD is collected in transmission mode using a FReLoN camera with fiber-optic taper coupling to the scintillator. The nano-XAS data are collected in XRF mode using a Si(111) double-crystal monochromator. Distribution maps and normalized line profiles collected along the PR heterostructure from XRF, XEOL, and XRD signals were processed with PyMca software. XRD data were integrated using Dioptas program.

**Electron Characterization:** The morphology and structure of the obtained PR heterostructure were analyzed by SEM, with a LEO-1530 SEM equipment, and by HRTEM, using a JEOL 2100 microscope working at 200 kV. The cross-sectional samples were obtained by means of a JEOL 4500 focused-ion beam/SEM with a 30 kV ion column and a Ga⁺ ion source.

**Supporting Information**

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

**Acknowledgements**

The authors acknowledge the ESRF for the beamtime allocated on the beamline ID16B (experiment: MA-4877). The authors thank Ana M. Sánchez (University of Warwick, UK), for their assistance with the TEM images acquisition. This work was financially supported by MICINN projects RTI2018-097195-B-100 and PID2021-122562NB-100. G.M.-C and F.M. acknowledge the financial support by European Union's Horizon 2020 research and innovation program under the Marie Skłodowska-Curie grant agreement n° 956548, Quantimony ITN Network.

**Conflict of Interest**

The authors declare no conflict of interest.

**Author Contributions**

G.M.-C. conceived the experiment; J.D., F.M., J.S., M.T., B.M. and G.M.-C. performed the synchrotron measurements; I.S. acquired the SEM images and R.M. carried out the DFT calculations; J.D. performed the data analysis; J.D. and G.M.C wrote the manuscript. All authors were involved in the discussion and commented on the manuscript.

**Data Availability Statement**

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

**Keywords**

nanowire heterostructures, oxide semiconductors, Plateau–Rayleigh instability, transparent conducting oxides, X-ray nanoprobes

Received: February 20, 2023

Revised: May 2, 2023

Published online: May 26, 2023

[1] L. J. Lauhon, M. S. Gudiksen, D. Wang, C. M. Lieber, *Nature* **2002**, *420*, 57.

[2] S. Mukherjee, U. Givan, S. Senz, A. Bergeron, S. Francoeur, M. De La Mata, J. Arbiol, T. Sekiguchi, K. Itoh, D. Isheim, D. N. Seidman, *Nano Lett.* **2015**, *15*, 3885.

[3] M. Heiss, Y. Fontana, A. Gustafsson, G. Wüst, C. Magen, D. O'Regan, J. Luo, B. Ketterer, S. Conesa-Boj, A. Kuhlmann, J. Houel, *Nat. Mater.* **2013**, *12*, 439.

[4] R. F. Oulton, V. J. Sorger, T. Zentgraf, R.-M. Ma, C. Gladden, L. Dai, G. Bartal, X. Zhang, *Nature* **2009**, *461*, 629.

[5] S. Nadj-Perge, S. Frolov, E. Bakkers, L. P. Kouwenhoven, *Nature* **2010**, *468*, 1084.

[6] J. Lahnemann, M. Den Hertog, P. Hille, M. de la Mata, T. Fournier, J. Schormann, J. Arbiol, M. Eickhoff, E. Monroy, *Nano Lett.* **2016**, *16*, 3260.

[7] P. Krogstrup, H. I. Jørgensen, M. Heiss, O. Demichel, J. V. Holm, M. Aagesen, J. Nygard, A. Fontcuberta i Morral, *Nat. Photonics* **2013**, *7*, 306.

[8] C. Jia, Z. Lin, Y. Huang, X. Duan, *Chem. Rev.* **2019**, *119*, 9074.

[9] G.-S. Liu, T. Wang, Y. Wang, H. Zheng, Y. Chen, Z. Zeng, L. Chen, Y. Chen, B.-R. Yang, Y. Luo, Z. Chen, *Nano Res.* **2022**, *15*, 2582.

[10] G.-S. Liu, M. He, T. Wang, L. Wang, Z. He, R. Zhan, L. Chen, Y. Chen, B.-R. Yang, Y. Luo, Z. Chen, *ACS Appl. Mater. Interfaces* **2020**, *12*, 53984.

[11] R. W. Day, M. N. Mankin, R. Gao, Y.-S. No, S.-K. Kim, D. C. Bell, H.-G. Park, C. M. Lieber, *Nat. Nanotechnol.* **2015**, *10*, 345.

[12] E. Sutter, P. Sutter, *Chem. Mater.* **2019**, *31*, 8174.

[13] J. Shi, J. Zhang, L. Yang, M. Qu, D.-C. Qi, K. H. Zhang, *Adv. Mater.* **2021**, *33*, 2006230.

[14] Z. Wu, W. Jie, Z. Yang, J. Hao, *Mater. Today Nano* **2020**, *12*, 100092.

[15] K. Wang, D. Liu, P. Deng, L. Liu, S. Lu, Z. Sun, Y. Ma, Y. Wang, M. Li, B. Y. Xia, C. Xiaoa, S. Ding, *Nano Energy* **2019**, *64*, 103954.

[16] S. A. Dayeh, J. Wang, N. Li, J. Y. Huang, A. V. Gin, S. T. Picraux, *Nano Lett.* **2011**, *11*, 4200.

[17] M. Huijben, G. Koster, M. K. Kruize, S. Wenderich, J. Verbeeck, S. Bals, E. Slooten, B. Shi, H. J. Molegraaf, J. E. Kleibeuker, S. van Aert, J. B. Goedkoop, A. Brinkman, D. H. A. Blank, M. S. Golden, G. van Tendeloo, H. Hilgenkamp, G. Rijnders, *Adv. Funct. Mater.* **2013**, *23*, 5240.

[18] Y. Li, L. Huang, B. Li, X. Wang, Z. Zhou, J. Li, Z. Wei, *ACS Nano* **2016**, *10*, 8938.

[19] S. Martín Sánchez, M. Botifoll, E. Oksenberg, C. Koch, C. Borja, M. C. Spadaro, V. Di Giulio, Q. Ramasse, F. J. García de Abajo, E. Joselevich, J. Arbiol, *Nat. Commun.* **2022**, *13*, 4089.

[20] T. Walbert, F. Muench, Y. Yang, U. Kunz, B.-X. Xu, W. Ensinger, L. Molina-Luna, *ACS Nano* **2020**, *14*, 11309.

[21] Z. Liu, D. Elbert, C.-L. Chien, P. C. Searson, *Nano Lett.* **2008**, *8*, 2166.

[22] S. K. Lim, M. Brewster, F. Qian, Y. Li, C. M. Lieber, S. Gradecak, *Nano Lett.* **2009**, *9*, 3940.

[23] H. Gao, W. Sun, Q. Sun, H. H. Tan, C. Jagadish, J. Zou, *Nano Lett.* **2019**, *19*, 3782.
[24] M. Hetzl, M. Kraut, J. Winnerl, L. Francaviglia, M. Düoblinger, S. Matich, A. Fontcuberta i Morral, M. Stutzmann, Nano Lett. 2016, 16, 7098.

[25] R. Songmuang, L. T. T. Giang, J. Bleuse, M. Den Hertog, Y.-M. Niquet, L. S. Dang, H. Mariette, Nano Lett. 2016, 16, 3426.

[26] S. Xu, P. Li, Y. Lu, Nano Res. 2018, 11, 625.

[27] R. W. Day, M. N. Mankin, C. M. Lieber, Nano Lett. 2016, 16, 2830.

[28] W. Li, Q. Jia, X. Liu, J. Zhang, Ceram. Int. 2017, 43, 2950.

[29] Y. Chu, S. Jing, X. Yu, Y. Zhao, Cryst. Growth Des. 2018, 18, 2941.

[30] Z. Xue, M. Xu, Y. Zhao, J. Wang, X. Jiang, L. Yu, J. Wang, J. Xu, Y. Shi, K. Chen, P. R. I. Cabarrocas, Nat. Commun. 2016, 7, 12836.

[31] K. M. Ridings, S. C. Hendy, Sci. Rep. 2022, 12, 20052.

[32] V. N. Gorshkov, V. V. Tereshchuk, P. Sareh, CrystEngComm 2020, 22, 2601.

[33] M. Garín, C. Jin, D. Cardador, T. Trifonov, R. Alcubilla, Sci. Rep. 2017, 7, 7233.

[34] R. H. Allaire, L. Kondic, L. J. Cummings, P. D. Rack, M. Fuentes-Cabrera, J. Phys. Chem. C 2021, 125, 5723.

[35] Z. Xue, W. Chen, X. Meng, J. Xu, Y. Shi, K. Chen, L. Yu, P. R. I. Cabarrocas, Appl. Surf. Sci. 2022, 593, 153435.

[36] Y. Li, T.-T. Zhuang, C. Zhang, L. Wu, S.-K. Han, S.-H. Yu, Acc. Mater. Res. 2020, 1, 126.

[37] G.-S. Liu, H. Zheng, Z. Zeng, Y. Wang, W. Guo, T. Wang, H. Chen, Y. Chen, S. Hu, L. Chen, Y. Chen, Nano Res. 2022, 15, 4552.

[38] J. Dolado, K. L. Renforth, J. E. Nunn, S. A. Hindsmarsh, P. Hidalgo, A. M. Sánchez, B. Méndez, Cryst. Growth Des. 2019, 20, 506.

[39] G. Martínez-Criado, J. Villanova, R. Tucoulou, D. Salomon, J.-P. Suuronen, S. Labouré, C. Guilloud, V. Valls, R. Barrett, E. Gagliardini, Y. Dabin, J. Synchrotron Radiat. 2016, 23, 344.

[40] D. Maestre, A. Cremades, J. Piqueras, J. Appl. Phys. 2004, 95, 3027.

[41] D. R. Miller, R. E. Williams, S. A. Akbar, P. A. Morris, D. W. McComb, Sens. Actuators, B: Chem. 2017, 240, 193.

[42] X. Zhou, F. Heigl, M. Murphy, T. Sham, T. Regier, I. Coulthard, R. Blyth, Appl. Phys. Lett. 2006, 89, 213109.

[43] X. Zhou, J. Zhou, M. Murphy, J. Ko, F. Heigl, T. Regier, R. Blyth, T. Sham, J. Chem. Phys. 2008, 128, 144703.

[44] J. Dolado, P. Hidalgo, B. Méndez, Mater. Lett. 2020, 275, 128152.

[45] A. Kar, S. Kundu, A. Patra, J. Phys. Chem. C 2011, 115, 118.

[46] X. Yuan, L. Lazzarini, G. Salviati, M. Zha, T. Sekiguchi, Mater. Sci. Semicond. 2006, 9, 331.

[47] M. A. Reshchikov, H. Morkoç, J. Appl. Phys. 2005, 97, 5.

[48] S. C. Cruz, S. Keller, T. E. Mates, U. K. Mishra, S. P. DenBaars, J. Cryst. Growth 2009, 311, 3817.

[49] N. Han, A. T. Hui, F. Wang, J. J. Hou, F. Xiu, T. Hung, J. C. Ho, Appl. Phys. Lett. 2011, 99, 083114.

[50] G. Martínez-Criado, A. Homs, B. Alén, J. A. Sans, J. Segura-Ruiz, A. Molina-Sánchez, J. Susini, J. Yoo, G.-C. Yi, Nano Lett. 2012, 12, 5829.

[51] J. Haug, A. Chassé, M. Dubiel, C. Eisenschmidt, M. Khalid, P. Esquinazi, J. Appl. Phys. 2011, 110, 063507.

[52] X. Liu, J. Iqbal, Z. Wu, B. He, R. Yu, J. Phys. Chem. C 2010, 114, 4790.

[53] D. Grandjean, R. E. Benfield, C. Nayral, A. Maisonnat, B. Chaudret, J. Phys. Chem. B 2004, 108, 8876.

[54] Q. Zhu, Q. Ma, D. Buchholz, R. Chang, M. J. Bedzyk, T. O. Mason, J. Appl. Phys. 2014, 115, 033512.