# Enhanced Optoelectronic Properties of Silicon Clathrate Films via Sodium Extraction by Vacuum Annealing

Anil Kumar Bharwal, Joseph P Briggs, Charif Tamin, Maxime Hanauer, Marc

Lenertz, Jérémy Bartringer, Stéphane Roques, Céline Chevalier, Mathieu

Gallart, Pierre Gilliot, et al.

► To cite this version:

Anil Kumar Bharwal, Joseph P Briggs, Charif Tamin, Maxime Hanauer, Marc Lenertz, et al. "Enhanced Optoelectronic Properties of Silicon Clathrate Films via Sodium Extraction by Vacuum Annealing." physica status solidi (RRL) - Rapid Research Letters (pss RRL), In press, 10.1002/pssr.202500318. [hal-05293689]

HAL Id: hal-05293689

https://hal.science/hal-05293689v1

Submitted on 2 Oct 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
# Enhanced Optoelectronic Properties of Silicon Clathrate Films via Sodium Extraction by Vacuum Annealing

Anil Kumar Bharwal,* Joseph P. Briggs, Charif Tamin, Maxime Hanauer, Marc Lenertz, Jérémy Bartringer, Stéphane Roques, Céline Chevalier, Mathieu Gallart, Pierre Gilliot, Marc Ziegler, Aziz Dinia, Abdelilah Slaoui, Reuben T. Collins, and Thomas Fix

Silicon clathrates, which exhibit a direct or near-direct bandgap, represent a promising alternative to conventional diamond cubic silicon for optoelectronic applications. They combine the elemental abundance and nontoxicity of Si with a more favorable bandgap and higher optical absorption at room temperature. Their open cage-like frameworks allow for reversible guest atom incorporation with minimal lattice distortion, making them attractive not only for optoelectronics but also for energy storage applications. However, conventional synthesis methods typically require the use of sodium as a structure-directing agent to stabilize the metastable clathrate lattice during thermal decomposition. The resulting sodium incorporation degrades the optoelectronic properties and complicates device integration. Herein, a nontoxic, scalable, vacuum-annealing approach is utilized that effectively removes sodium from silicon clathrates films without the use of hazardous chemicals, such as iodine, HF, or $SF_6$. Optoelectronic characterization reveals a significant improvement in semi-conducting behavior, including enhanced surface photovoltage response, marking a critical step toward developing sodium-free silicon clathrate thin films as efficient light absorbers. This vacuum annealing process is compatible with large-area substrates and offers a key step toward a safer, and more sustainable approach to integrating exotic Si phases into the next-generation optoelectronic devices.

## 1. Introduction

Silicon plays a central role in modern electronics and solar energy technologies due to its abundance, thermal stability, doping versatility, and ability to form stable surface oxides.[1] Similar to carbon and sulfur, silicon can exist in multiple crystalline forms, some of which may exhibit direct bandgaps and strong light absorption.[2] Such silicon allotropes have emerged as promising candidates for low-cost, high-efficiency thin-film solar cells that remain compatible with existing silicon processing technologies.[3,4]

In the early 1960s, researchers at the University of Bordeaux in France began experimentally investigating silicon clathrates, proposing that they form guest-stabilized cage-like structures, analogous to gas hydrates.[5,6] The first experimental confirmation of these low-density silicon frameworks came in 1965, when Kasper et al. identified silicon clathrates in powdered form.[7] This discovery demonstrated that silicon can adopt expanded cage-like frameworks capable of hosting various guest atoms.[8–14] Since then, continual

improvements have been made in the characterization of their structural and compositional features.[15] These cage-like

A. K. Bharwal, M. Hanauer, J. Bartringer, S. Roques, A. Slaoui, T. Fix ICube Laboratory
University of Strasbourg and CNRS
67037 Strasbourg, France
E-mail: bharwal@unistra.fr

J. P. Briggs
Department of Chemical and Biological Engineering
Colorado School of Mines, Golden
Colorado 80401, USA

<figure><img src="image_2.png" /></figure>

The ORCID identification number(s) for the author(s) of this article can be found under https://doi.org/10.1002/pssr.202500318.

© 2025 The Author(s). physica status solidi (RRL) Rapid Research Letters 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.

DOI: 10.1002/pssr.202500318

C. Tamin, C. Chevalier
CNRS
Ecole Centrale de Lyon
INSA Lyon
Université Claude Bernard Lyon 1
CPE Lyon
INL
UMR5270
69100 Villeurbanne, France

M. Hanauer, M. Lenertz, M. Gallart, P. Gilliot, M. Ziegler, A. Dinia
IPCMS Laboratory
University of Strasbourg and CNRS
67037 Strasbourg, France

R. T. Collins
Department of Physics
Colorado School of Mines, Golden
Colorado 80401, USA
structures are distinct from the diamond cubic silicon (d-Si) tetrahedral network and have attracted growing interest due to their tunable structural and electronic properties.[16] Si clathrates have a direct or near-direct bandgap and exhibit enhanced absorption in the visible spectrum compared to conventional d-Si, which also has an indirect bandgap of ≈1.1 eV.[17,18]

Silicon clathrates primarily exist as two structural forms: type-I (MgSi₄₆) and type-II (MₓSi₁₃₆). Type-I has fully occupied cages and exhibits metallic behavior, with potential applications in thermoelectrics.[5,19] However, the removal of the stabilizing guest atom, typically sodium, has not yet been demonstrated, which limits the potential of Type-I for optoelectronic applications. In contrast, type-II clathrates feature larger polyhedral cages characterized by more open geometries and expanded cavity volumes, allowing for greater flexibility in sodium occupancy.[20,21] This structural adaptability permits the tuning of the optoelectronic properties by adjusting the sodium content (0 < x < 24).[22]

When the sodium level falls below the threshold for the metal-insulator transition, type-II clathrates are expected to shift from metallic to n-type semiconducting behavior, with predicted optical bandgaps greater than 1.9 eV.[23,24] Type-II clathrates have been explored for applications in photovoltaics,[25-28] batteries,[11,29-31] spintronics,[32] and catalysis.[33] Their wide bandgaps make them particularly well suited for indoor photovoltaic (IPV) devices, where absorption in the 400-700 nm range aligns with the emission spectra of common indoor lighting sources such as LEDs and fluorescent lamps.[34,35]

Given the potential of type-II Si clathrates, it is essential to optimize their optoelectronic properties by reducing the sodium content. However, traditional methods often rely on chemical agents such as iodine, SF₆, or other complex procedures.[36-39] These approaches are hazardous, expensive, and require regulatory authorization. Others have reported a reduction in Na concentration in Si clathrate powders with prolonged annealing under vacuum.[40] Vacuum annealing offers a more refined and controlled alternative for achieving this goal.

Following this approach, we present a safe, scalable, and nontoxic method for decreasing the sodium content in silicon clathrate films using extended vacuum annealing. Using two distinct vacuum conditions, we showed that extended high-vacuum treatment effectively removes sodium, even from large-area films, without the use of hazardous chemicals. Beyond sodium removal, this process enhances the optoelectronic properties of the films, increasing the bandgap from 1.70 up to 1.74 eV, which is well within the ideal range (1.6-2.0 eV) for photovoltaic applications.[41] The widened bandgap improves the compatibility with tandem solar cell architectures, making sodium-free clathrates promising absorber materials for both single- and multijunction photovoltaic devices. Among the methods investigated, vacuum annealing emerges as the most practical and environment-friendly method for producing high-quality sodium-reduced clathrate films. However, deeper insights into postannealing energy levels and electronic structures are required to fully optimize these clathrate films for optoelectronic applications.

To address this, a detailed understanding of the energy levels of the material is essential for selecting suitable contact materials with matching work function (WF) in high-performance

optoelectronic devices.[42] In this study, we present an energy-band diagram using an integrated system comprising a Kelvin probe (KP), surface photovoltage (SPV), ambient-pressure photo-emission spectroscopy (APS), and optical measurements.[43-45] A strong SPV response is often indicative of good photovoltaic behavior. Here SPV signal on optimized films was 165 mV, one of the largest value ever obtained so far, highlighting the potential of sodium-free clathrate films for device applications. Overall, sodium was successfully removed from the clathrate cages via vacuum annealing, without the use of hazardous chemicals, and the resulting films exhibited enhanced optoelectronic properties.

## 2. Results and Discussion

Once clathrate films were fabricated, they were subjected to vacuum annealing under various conditions. To evaluate the impact of annealing on the film properties, a comprehensive set of structural, morphological, optical, and electrical characterizations were performed.

### 2.1. Structural and Electronic Characterization

X-ray diffraction (XRD) measurements were performed at room temperature to confirm the formation of the silicon clathrate phase and to assess the effects of sequential vacuum annealing on crystallinity and sodium content. **Figure 1a** shows the θ-2θ XRD patterns of the as-grown clathrate films before and after low- and high-vacuum annealing. The as-grown film shows the characteristics peaks of the type-II NaₓSi₁₃₆ phase (marked by red star ⋆), indicating a high sodium content, consistent with previous studies.[15] All peaks corresponding to ICDD sheet number 00-018-1244 are present, along with minor secondary phases, including peaks corresponding to type-I Na₈Si₄₆ (blue triangles ▲) and crystalline c-Si (orange diamonds ♦). The intensity ratio of the peaks around 2θ ≈20° ((311) and (222) notably) relative to the

<figure><img src="image_3.png" /><figcaption>**Figure 1.** a) Room temperature θ-2θ XRD patterns of silicon clathrate films: as-grown, low-vacuum annealed for 2 days (LV-2D), and high-vacuum annealed for 3 days (HV-3D). b) Thermally pressed film before and after 6 days of high-vacuum annealing (pressed-6D). The bottom panels show reference diffraction patterns for both phases to assist in peak identification.</figcaption></figure>
(822) peak around 53° is notably lower in the as-grown and low-vacuum annealed samples, consistent with higher sodium occupancy, as previously reported in the literature.[22] High-vacuum annealing led to a relative increase in the intensity of the 20° peaks, indicating reduced sodium content while preserving the clathrate structure. These peaks are known to be inversely related to sodium content, and the observed trend aligns well with earlier findings that associate their enhancement with progressive sodium extraction.[36,46,47] Building on previous research demonstrating the suitability of pressed clathrate films for device integration,[46] we examined the influence of vacuum annealing on the pressed films. The pressed samples were subjected to extended high-vacuum annealing for 6 days. As shown in Figure 1b, the sodium-rich phases were diminished compared to those of the unannealed pressed film. Both samples retained the primary type-II clathrate structure. These observations were further confirmed by Rietveld analysis, which provided a semi-quantitative assessment of the phase composition and sodium site occupancies, allowing us to estimate the sodium content in the samples. As shown in Figure S1, Supporting Information, the type-II phase accounted for >90 at.%, while a minor presence of the type-I clathrate is consistently detected. Upon high-vacuum processing (HV-3D and pressed-6D), the type-I phase was further reduced, accompanied by the formation of cubic Si. Refinement of the sodium site occupancies confirmed sodium extraction from the type-II phase. Additionally, the refinement agreement factors (χ², Rₚ, R<sub>wp</sub>, and R<sub>bragg</sub>) for each sample are listed in Table S1, Supporting Information, confirming the reliability of the fits.

The surface morphology and thicknesses of the clathrate films
were analyzed using scanning electron microscope (SEM), as

shown in **Figure 2**. The as-grown film (Figure 2a) exhibits a sur-
face composed of tightly packed clathrate aggregates. Annealing
at 400 °C under low-vacuum conditions resulted in a pronounced
increase in surface roughness (Figure 2b). Subsequent annealing
under high-vacuum conditions caused an additional, though
less pronounced, increase in the surface roughness (Figure 2c),
indicating that vacuum annealing gradually affected the film
morphology. Thermal pressing had a notable impact on the mor-
phology, resulting in a much smoother surface (Figure 2a,d)
and a decrease in the film thickness by approximately by half,
as shown in Figure S2, Supporting Information. In addition,
Figure 2e illustrates that prolonged high-vacuum annealing
introduces noticeable roughness and small cracks on the surface,
likely due to thermal stress and extended sodium extraction,
which can lead to microstructural rearrangements and changes
in the grain boundaries.

Elemental composition analysis was performed using energy-
dispersive X-ray spectrometry (EDS), with particular focus on
sodium concentration (%). Although EDS has limitations in
accurately quantifying light elements, such as Na, it remains
useful for tracking relative trends across successive vacuum-
annealing steps. As illustrated in Figures S3 and S4, Supporting
Information, a gradual decrease in sodium concentration was
observed with increasing annealing duration and vacuum level.
Figure 2f presents a comparison between these values and those
derived from XRD Rietveld refinement, illustrating a notable rel-
ative decrease in sodium content, particularly following high-
vacuum treatment. Thus, prolonged high-vacuum annealing is
an effective method for extracting sodium from the clathrate cages.

Raman spectroscopy was used to complement the XRD
analysis and to confirm the presence and evolution of the

<figure><img src="image_4.png" /><figcaption>**Figure 2.** Influence of thermal treatment on the morphology of Si clathrate films. SEM images of clathrate films: a) as-grown film, displaying a rough and cracked surface; b) postannealing at 400 °C in a low-vacuum environment, showing increased surface roughness; c) same sample after extended high-vacuum annealing for 3 days; d) as-grown film that has been thermally pressed; e) thermally pressed film after 6 days of high-vacuum annealing; and f) sodium ratio for each sample measured by XRD and EDS.</figcaption></figure>
NaₓSi₁₃₆ phase under different vacuum-annealing conditions (Figure S5a–b, Supporting Information). The Raman spectra for both the as-grown and vacuum-annealed samples (in low and high vacuum) displayed vibrational modes characteristic of type-II NaₓSi₁₃₆, consistent with previous findings.[48] As shown in Figure S5a, Supporting Information, the Raman features of the as-grown sample are broad and poorly defined, suggesting structural disorder. After vacuum annealing, the full width at half maximum (FWHM) values decrease, indicating improved crystallinity (Figure S5c, Supporting Information). Thermal pressing of the as-grown sample enhanced the film density and resulted in sharper and more intense Raman peaks, owing to the reduced surface roughness and better structural order (Figure S5b, Supporting Information). Prolonged high-vacuum annealing of the pressed sample over 6 days led to narrower peaks and reduced FWHM values, which is consistent with the progressive sodium removal and improved crystallinity observed in XRD.

To further examine the film quality before and after annealing, electron paramagnetic resonance (EPR) was performed on a representative as-grown and pressed samples vacuum annealed for 6 days. The type-II clathrate EPR with sodium guest has some interesting features detailed in prior studies.[36,49,50] In EPR, the derivative of the absorption spectrum is measured, as shown in **Figure 3a**, which provides more resolution to features than the absorption peak (Figure 3b), which was determined by integrating the spectra. Sharp features are shown in Figure 3a originate from isolated donor electrons interacting with the spin 3/2 Na nucleus, resulting in the four hyperfine lines occurring at 311, 324, 337, and 350 mT at this microwave frequency. Additionally, there were small features between the hyperfine lines associated with two Na atoms in neighboring cages observed at 318 and 344 mT, with the other features obscured by the hyperfine lines. The broad peak centered at 330 mT, which can be better seen in the absorption spectra in Figure 3b, originates from multiple Na-forming clusters with absorption peaks overlapping from various numbers of Na in neighboring cages.[36,49] These three features change in relative intensities between the as-grown (black) and pressed-6D (purple) films, notably, the large cluster feature diminished significantly, corroborate the sodium removal trend observed in the XRD analysis. Additionally, the two Na lines and isolated Na hyperfine lines grew after annealing and pressing, indicating the transition of large clusters of Na to more isolated Na spread out between cages. This further suggests sodium content reduction, as proposed by XRD. An additional noticeable

<figure><img src="image_5.png" /><figcaption>**Figure 3.** a) Room-temperature EPR spectra of the as-grown and vacuum-annealed (pressed-6D) silicon clathrate films, showing distinct paramagnetic defect features. b) Low-temperature integrated EPR absorption spectra indicating changes in defect density upon annealing.</figcaption></figure>

change was observed at 335 mT and associated with untermi-
nated silicon bonds, or dangling bonds.[49] These dangling bond
defects are shown to reduce significantly between the as-grown
and pressed-6D samples becoming just a small bump to the left
of the third hyperfine line in Figure 3a. This could be due to the
restoration of lattice defects by the pressing or vacuum-annealing
procedures. While XRD has been shown to be limited in its
ability to accurately estimate sodium content below $x \approx 1$,[11,36]
EPR spin density analysis provides a more sensitive estimate
using the methodology described in Ref. [49], spin counting
revealed a sodium spin density before treatment of $\approx 1 \times 10^{20}$
Na carriers/cm$^3$ ($x \approx 0.3$) and after 6 days of vacuum annealing,
the spin density dropped to $\approx 5 \times 10^{18}$ carriers/cm$^3$ correspond-
ing to $x \approx 0.02$. This substantial reduction confirms the effective-
ness of vacuum annealing in lowering sodium content to levels
suitable for semiconducting applications.

## 2.2. Optoelectronic Characterization

Having established the postannealing structural integrity, we then
evaluated the impact of these modifications on the optoelectronic
behavior. Photoluminescence (PL) measurements were conducted
using an excitation wavelength of 532 nm to assess the bandgap
changes resulting from the various annealing processes. The
optical bandgap of each sample was determined from the peak
positions of the PL emission spectra. The shifts seen in the peak
of samples with differing Na concentrations and treatments would
be the same shifts in the leading edge of the PL peak and so can be
associated with changes in the bandgap. As shown in **Figure 4a**,
the bandgap of the as-grown film was ≈1.70 eV. After annealing
for 2 days in low vacuum, the bandgap remained unchanged.
When the same sample (LV-2D) was further subjected to three
additional days annealing at 400 °C in high vacuum, the PL peak
shifted to higher energy indicating a bandgap increased to 1.73 eV.
This shift supports our earlier observations and confirms that
the ongoing sodium extraction leads to a bandgap expansion.[16]
This increase is attributed to the enhanced structural order
and reduction in sodium-related defects and impurities.[38] As pre-
viously demonstrated, thermal pressing of the as-grown film
results in a denser structure and a more pronounced PL response
at 1.70 eV. Following extended 6 day high-vacuum annealing
of this pressed sample, the PL peak shifted further to ≈1.74 eV
(Figure 4b), indicating continued bandgap expansion and
improved optoelectronic quality. We observe a shoulder in the
PL spectra between 1.9 and 2.0 eV, present in all samples except
those subjected to high-vacuum annealing. This feature is
likely associated with defect-related states, and vanishes after high-
vacuum annealing indicates a decrease in surface or structural
defects. We recognize that factors such as roughness, inhomoge-
neity, and impurity defects may influence PL measurements.
Consequently, we rely on PL data from the type-II phase, which
yields consistent values aligned with previous reports[14,46]
(see Table S2, Supporting Information), and within the typical
uncertainty of optical measurements. However, the observed opti-
cal bandgap variation may not be statistically significant and could
be subject to debate due to surface effects.

After determining the bandgap evolution through steady-state
PL, we examined the carrier dynamics using time-resolved PL
(as-grown) to 5.68 eV (LV-2D), and to 5.37 eV (HV-3D), indicat-
ing a corresponding shift in the valence band (VB) edge. Thermal
pressing lowered the VB from 5.73 to 5.60 eV, and subsequent
extended vacuum annealing of the pressed film (pressed-6D)
reduced it further to 5.20 eV, as shown in Figure S6a and S7.
Notably, the measured IEs values were in good agreement with
our previously reported values for silicon clathrates.[16] Finally,
the conduction band (CB) positions were calculated by subtract-
ing the optical bandgap from the measured IE values. By com-
bining these measurements with the optical bandgap data
obtained from the PL, we built complete energy-band diagrams
for each sample, as shown in Figure 4f. The resulting energy-
band diagrams confirm that all processed films maintained their
n-type semiconducting nature after vacuum annealing. Figure
S8, Supporting Information illustrates the CB offset relative to
the Fermi level (CB-E<sub>f</sub>) for the different clathrate films. The
HV-3D sample showed the highest CB-E<sub>f</sub> value (0.86 eV), com-
pared to 0.67 eV for both the as-grown and LV-2D samples. This
suggests that high-vacuum annealing pushes the Fermi level fur-
ther into the bandgap, thereby enhancing semiconducting prop-
erties. This is consistent with a sodium content reduction as
observed in XRD, SEM, and EPR. Conversely, the pressed sam-
ple experiences a slight reduction to 0.65 eV, likely due to struc-
tural or electronic changes during thermal processing, but this
value rises to 0.84 eV after prolonged annealing under high vac-
uum (pressed-6D). Overall, these results confirm that annealing
under high-vacuum conditions significantly improves the semi-
conducting characteristics of Na<sub>x</sub>Si<sub>1-3x</sub> clathrate films, highlight-
ing their ability to adjust energy levels and their suitability for
incorporation into various optoelectronic device designs.

2.3. SPV

The SPV, which is characterized by the variation in contact poten-
tial difference (CPD) between dark and illuminated conditions,
serves as a crucial measure of photoresponse. We used KP meas-
urements under white-light illumination to evaluate the CPD of
clathrate films, which is particularly effective for assessing photo-
voltaic properties.[43] **Figure 5a** shows the transient-SPV behavior
under alternating dark and illuminated conditions, while
Figure 5b displays the corresponding SPV signals as a function
of light intensity. The as-grown film exhibited a weak SPV signal,
because of its high sodium content, which may promote charge
trapping and inhibit efficient carrier separation.[16] Following
annealing at 400 °C in low vacuum, the SPV increased slightly,
indicating slight sodium extraction. Subsequent high-vacuum
annealing led to a more significant improvement, suggesting
reduced recombination and improved electronic quality. To meet
the demands of photovoltaic integration, it is crucial to fabricate
thin and compact films that facilitate efficient carrier transport.[46]
Although thermally pressed clathrate films exhibited only slight
improvements compared to the as-grown film, an extended 6 days
annealing process in a high vacuum resulted in a notable increase
in the photoresponse, with the maximum SPV reaching ≈165 mV
(Figure 5c,d). Beyond its role as a photoresponse metric, SPV also
serves as an indirect indicator of the minority carrier diffusion
length,[53] and a stable response across different light intensities
is advantageous for photovoltaic applications.[44,54] These findings

<figure><img src="image_6.png" /><figcaption>**Figure 5.** a,c) Time-resolved CPD measurements of clathrate films under alternating dark and halogen visible-light illumination, obtained using KP. b,d) Corresponding SPV responses recorded under varying halogen light intensities.</figcaption></figure>

emphasize the combined benefits of sodium extraction, structural
densification, and extended vacuum treatment in enhancing the
optoelectronic quality of the NaₓSi₁₃₆ films. This highlights the
importance of engineering highly compact, Na-reduced clathrate
structures to optimize the photovoltaic performance.

**3. Conclusion**

In this study, we demonstrated that vacuum annealing offers a
clean, scalable, and effective approach for adjusting the optoelec-
tronic properties of silicon clathrate films. Unlike traditional
chemical etching techniques, that rely on dangerous reagents,
such as SF₆ or iodine, vacuum annealing enables controlled
sodium extraction in an eco-friendly manner. Notably, SPV meas-
urements showed a consistent enhancement in photoresponse
with progressive sodium depletion and reduced type-I phase
impurities, attributed to improved carrier dynamics and lower
recombination losses. By integrating the KP, APS, and PL data,
we constructed comprehensive, detailed energy-band diagrams
that illustrate the shifts in the Fermi level, CBs and VBs, and
WF alignment. These insights provide a valuable foundation
for optimizing the band alignment in clathrate-based heterojunc-
tions and advancing their integration into next-generation solar
cells. Future studies will explore the economic feasibility and
scalability of this annealing method. Overall, our results establish
vacuum annealing as a sustainable and versatile method for engi-
neering the electronic structures of clathrate semiconductors, pav-
ing the way for their use in high-performance optoelectronic and
photovoltaic devices, while also aligning with recent findings on
hydrogen incorporation effects.[55]

**4. Experimental Section**

**Materials**: P-type silicon wafers (525 ± 25 µm thickness, 1–5 Ω cm resistivity) were purchased from BT Electronics. Sodium metal (99.8% purity),
[14] A. K. Bharwal, M. Hanauer, C. Tamin, R. Vollondat, S. Roques, J. Bartringer, D. Stoeffler, C. Chevalier, A. Dinia, A. Slaoui, T. Fix, *MRS Adv.* **2024**, 9, 1685.

[15] R. Vollondat, S. Roques, C. Chevalier, J. Bartringer, J.-L. Rehspringer, A. Slaoui, T. Fix, *J. Alloys Compd.* **2022**, *903*, 163967.

[16] A. K. Bharwal, R. Vollondat, C. Tamin, S. Roques, J. Bartringer, D. Stoeffler, C. Chevalier, A. Dinia, A. Slaoui, T. Fix, *ACS Appl. Energy Mater.* **2024**, *7*, 8554.

[17] M. A. Green, M. J. Keevers, *Prog. Photovolt. Res. Appl.* **1995**, *3*, 189.

[18] J. Gryko, P. F. McMillan, R. F. Marzke, G. K. Ramachandran, D. Patton, S. K. Deb, O. F. Sankey, *Phys. Rev. B* **2000**, *62*, R7707.

[19] J.-A. Dolyniuk, B. Owens-Baird, J. Wang, J. V. Zaikina, K. Kovnir, *Mater. Sci. Eng. R Rep.* **2016**, *108*, 1.

[20] J. G. Slingsby, N. A. Rorrer, L. Krishna, E. S. Toberer, C. A. Koh, C. M. Maupin, *Phys. Chem. Chem. Phys.* **2016**, *18*, 5121.

[21] V. I. Smelyansky, S. T. John, *Chem. Phys. Lett.* **1997**, *264*, 459.

[22] S. Stefanoski, C. D. Malliakas, M. G. Kanatzidis, G. S. Nolas, *Inorg. Chem.* **2012**, *51*, 8686.

[23] S. B. Roy, K. E. Sim, A. D. Caplin, *Philos. Mag. B* **1992**, *65*, 1445.

[24] J. C. Conesa, C. Tablero, P. Wahnon, *J. Chem. Phys.* **2004**, *120*, 6142.

[25] J. Wu, H. Gao, K. Xia, D. Xing, J. Sun, *Appl. Phys. Lett.* **2017**, *111*, 173904.

[26] N. A. Mahammedi, H. Gueffaf, B. Lagoun, M. Ferhat, *Opt. Mater.* **2020**, *107*, 110043.

[27] T. Fix, R. Vollondat, A. Ameur, S. Roques, J.-L. Rehspringer, C. Chevalier, D. Muller, A. Slaoui, *J. Phys. Chem. C* **2020**, *124*, 14972.

[28] C. Tamin, A. K. Bharwal, C. Chevalier, A. Fave, E. Fourmond, S. Roques, A. Dinia, A. Slaoui, T. Fix, *Adv. Mater. Interfaces* **2025**, *12*, e00301.

[29] J. Yang, S. T. John, *J. Mater. Chem. A* **2013**, *1*, 7782.

[30] P. Warrier, C. A. Koh, *Appl. Phys. Rev.* **2016**, *3*, 040805.

[31] A. Dopilka, J. M. Weller, A. Ovchinnikov, A. Childs, S. Bobev, X. Peng, C. K. Chan, *Adv. Energy Sustainable Res.* **2021**, *2*, 2000114.

[32] J. P. Briggs, Y. Liu, P. C. Taylor, M. Singh, R. T. Collins, C. A. Koh, *Appl. Phys. Lett.* **2024**, *125*, 202402.

[33] W. Al Maksoud, R. U. S. Bacha, J. Pixius, M. Viswanathan, Y. Vaishnav, R. K. Rai, M. N. Hedhili, Q. Wang, Y. Kobayashi, *Adv. Mater.* **2024**, *36*, 2406944.

[34] A. Chakraborty, G. Lucarelli, J. Xu, Z. Skafi, S. Castro-Hermosa, A. B. Kaveramma, R. G. Balakrishna, T. M. Brown, *Nano Energy* **2024**, *128*, 109932.

[35] D. Müller, E. Jiang, P. Rivas-Lazaro, C. Baretzky, G. Loukeris, S. Bogati, S. Paetel, S. J. C. Irvine, O. Oklobia, S. Jones, D. Lamb, A. Richter, G. Siefer, D. Lackner, H. Helmers, C. Teixeira, D. Forgács, M. Freitag, D. Bradford, Z. Shen, B. Zimmermann, U. Würfel, *ACS Appl. Energy Mater.* **2023**, *6*, 10404.

[36] A. Ammar, C. Cros, M. Pouchard, N. Jaussaud, J.-M. Bassat, G. Villeneuve, M. Duttine, M. Ménétrier, E. Reny, *Solid State Sci.* **2004**, *6*, 393.

[37] T. Kume, F. Ohashi, K. Sakai, A. Fukuyama, M. Imai, H. Udono, T. Ban, H. Habuchi, H. Suzuki, T. Ikari, *Thin Solid Films* **2016**, *609*, 30.

[38] Y. Liu, W. K. Schenken, L. Krishna, A. A. A. Majid, T. E. Furtak, M. Walker, C. A. Koh, P. C. Taylor, R. T. Collins, *Appl. Phys. Rev.* **2021**, *8*, 041408.

[39] S. Iwasaki, H. Morito, T. Komine, K. Morita, T. Shibuya, J. Nishii, M. Fujioka, *Adv. Mater.* **2022**, *34*, 2106754.

[40] L. Krishna, L. L. Baranowski, A. D. Martinez, C. A. Koh, P. C. Taylor, A. C. Tamboli, E. S. Toberer, *CrystEngComm* **2014**, *16*, 3940.

[41] S. J. Adjogri, E. L. Meyer, *Materials* **2021**, *14*, 7857.

[42] L. Lin, R. Jacobs, T. Ma, D. Chen, J. Booske, D. Morgan, *Phys. Rev. Appl.* **2023**, *19*, 037001.

[43] J. R. Harwell, T. K. Baikie, I. D. Baikie, J. L. Payne, C. Ni, J. T. S. Irvine, G. A. Turnbull, I. D. W. Samuel, *Phys. Chem. Chem. Phys.* **2016**, *18*, 19738.

[44] C. Henderson, J. Luke, I. S. Bicalho, L. Correa, E. J. Yang, M. Rimmele, H. Demetriou, Y.-C. Chin, T. Lan, S. Heutz, *Energy Environ. Sci.* **2023**, *16*, 5891.

[45] Y. Fu, T. H. Lee, Y.-C. Chin, R. A. Pacalaj, C. Labanti, S. Y. Park, Y. Dong, H. W. Cho, J. Y. Kim, D. Minami, J. R. Durrant, J.-S. Kim, *Nat. Commun.* **2023**, *14*, 1870.

[46] A. K. Bharwal, J. P. Briggs, C. Tamin, M. Hanauer, R. Vollondat, J. Bartringer, S. Roques, C. Chevalier, A. Dinia, R. T. Collins, A. Slaoui, T. Fix, *ACS Appl. Energy Mater.* **2025**, *8*, 1752.

[47] R. Vollondat, D. Stoeffler, D. Preziosi, S. Roques, A. Slaoui, T. Fix, *J. Chem. Phys.* **2023**, *158*, 164709.

[48] O. I. Barkalov, M. A. Kuzovnikov, I. A. Sholin, N. S. Orlov, *Solid State Commun.* **2021**, *340*, 114492.

[49] W. K. Schenken, Y. Liu, L. Krishna, A. A. A. Majid, C. A. Koh, P. C. Taylor, R. T. Collins, *Phys. Rev. B* **2020**, *101*, 245204.

[50] H. Yahiro, K. Yamaji, M. Shiotani, S. Yamanaka, M. Ishikawa, *Chem. Phys. Lett.* **1995**, *246*, 167.

[51] S. M. Santos, B. Yuma, S. Berciaud, J. Shaver, M. Gallart, P. Gilliot, L. Cognet, B. Lounis, *Phys. Rev. Lett.* **2011**, *107*, 187401.

[52] A. Girault, M. Gallart, M. Ziegler, O. Crégut, M. Balestrieri, I. Gozhyk, H. Montigaud, D. Sciacca, L. Mancini, P. Gilliot, *Phys. Rev. B* **2023**, *108*, 155204.

[53] J. Lagowski, A. M. Kontkiewicz, L. Jastrzebski, P. Edelman, *Appl. Phys. Lett.* **1993**, *63*, 2902.

[54] J. Li, H. E. T. Mendoza, A. Kundmann, F. E. Osterloh, *Sustain. Energy Fuels* **2024**, *8*, 2494.

[55] M. Hanauer, A. K. Bharwal, Y. Tong, B. Aissa, C. Tamin, J. Bartringer, S. Roques, D. Muller, C. Chevalier, A. Dinia, A. Slaoui, T. Fix, *J. Phys. Energy* **2025**, *7*, 045019.

[56] N. Ha My Dang, P. Bouteyre, G. Trippé-Allard, C. Chevalier, E. Deleporte, E. Drouard, C. Seassal, H. S. Nguyen, *Opt. Mater. Express* **2024**, *14*, 1655.

[57] J. Rodríguez-Carvajal, *Phys. B Condens. Matter* **1993**, *192*, 55.

[58] S. Stoll, A. Schweiger, *J. Magn. Reson.* **2006**, *178*, 42.