# Extending mechanical size effect range of thin film metallic glasses by nanoengineering their atomic and nanostructure

Francesco Bignoli, Alex Lemarchand, Szilvia Kalácska, Dominique Thiaudiere, Philippe Djemia, Damien Faurie, Andrea Li Bassi, Matteo Ghidelli

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## ► To cite this version:

Francesco Bignoli, Alex Lemarchand, Szilvia Kalácska, Dominique Thiaudiere, Philippe Djemia, et al. Extending mechanical size effect range of thin film metallic glasses by nanoengineering their atomic and nanostructure. Acta Materialia, 2025, 300, pp.121456. 10.1016/j.actamat.2025.121456. hal-05231561

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https://hal.science/hal-05231561v1

Submitted on 3 Sep 2025

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Distributed under a Creative Commons Attribution 4.0 International License
Extending mechanical size effect range of thin film
metallic glasses by nanoengineering their atomic
and nanostructure

Francesco Bignoli¹,³*, Alex Lemarchand¹, Szilvia Kalácska², Dominique Thiaudiere⁴, Philippe Djemia¹, Damien

Faurie¹, Andrea Li Bassi³, Matteo Ghidelli¹*

¹Laboratoire des Sciences des Procédés et des Matériaux (LSPM), CNRS, Université Sorbonne Paris Nord, 93430, Villetaneuse, France.

² Mines Saint-Etienne, Univ Lyon, CNRS, UMR 5307 LGF, Centre SMS, 158 cours Fauriel 42023 Saint-Étienne, France.

³Dipartimento di Energia, Laboratorio Materiali Micro e Nanostrutturati, Politecnico di Milano, via Ponzio 34/3, I-20133 Milano, Italy.

⁴Soleil Synchrotron, French National Centre for Scientific Research, L'Orme des Merisiers Départementale 128, 91190 Saint-Aubin, France.

*Email: francescobignoli@cnr.it, matteo.ghidelli@lspm.cnrs.fr

**Abstract**

Nanoscale metallic glasses (MGs) have demonstrated unique strength – ductility balance with a yield strength close to the theoretical limit and large homogeneous plastic deformation, mitigating the shear band (SB) instability. However, such mechanical size effects are triggered only for intrinsic sample dimensions below few hundreds of nm, reducing their application range. In this work, we present a novel synthesis strategy to fabricate nanostructured thin film metallic glasses (TFMGs) with tailored local heterogeneities by exploiting the potential of Pulsed Laser Deposition (PLD), managing to extend MG's mechanical size effect range up to the micrometer scale. We realize this by synthesizing model Zr₅₀Cu₅₀ and Zr₄₆Cu₄₆Al₈ TFMGs with controlled
nanostructure compact and nanogranular, exploiting atom-by-atom and cluster-assembled growth regimes.

PLD-deposited TFMGs have unique atomic structures as observed by synchrotron X-ray scattering, with high mass density, resulting in enhanced elastic modulus (up to 120 GPa) and hardness (up to 8.6 GPa), significantly above sputter-deposited counterparts. Moreover, PLD-deposited TFMGs exhibit high yield strength ($\sigma_y$ >2.8 GPa) and up to 6.7% elastoplastic deformation which are maintained even at high strain rates up to 10 s⁻¹ for large micropillar dimension of 1.3 x 3 µm (diameter x height). Such values significantly surpass those of traditional families of bulk and TFMGs, where size effects on $\sigma_y$ and plasticity are activated for intrinsic dimensions < 0.5 µm. Overall, we show that PLD can produce nanostructured TFMGs with high strength and plasticity balance kept for large intrinsic dimensions, opening a new synthesis approach for strong and ductile materials.

Keywords: Pulsed Laser Deposition; Thin film metallic glasses; Extended plasticity; Mechanical properties; Local heterogeneities; Micropillar compression; Synchrotron measurements; Strain rate sensitivity.

# 1. Introduction

The fine control of the microstructure of metallic thin films is a necessary step in order to produce new nanoengineered materials with enhanced mechanical properties able to surpass the limits imposed by their base atomic structure. For example, metallic glasses (MGs) possess high yield strength close the theoretical atomic decohesion limit and elastic deformation above their crystalline counterparts [1, 2]. However, the absence of dislocations and grain boundaries lead to deformation occurring through nanometric (~10 nm) shear bands (SBs) which develop immediately after yielding and induce a macroscopically brittle behavior [3].

In order to mitigate this behavior, the first strategy consists in reducing the size of the sample below few hundreds of nanometers in order to trigger mechanical size effects, hindering the nucleation and propagation of SBs [4-6]. This has been demonstrated in the work of Tian et al. [4], who conducted in situ TEM tensile tests on 200-300 nm thick Cu₄₉Zr₅₁, as well as by Ghidelli et al. [5, 6], who performed microscale tensile test of ZrNi TFMGs with thickness <300 nm showing that yield strength ($\sigma_y$) can be enhanced up to the theoretical limit (3.3 GPa, atomic decohesion), together with a homogenous deformation (up to 15 %). Such geometrical
confinement can be further enhanced by combining different materials in nanolayered architectures, exploiting the presence of interfaces in order to block the SB propagation or adding intrinsically ductile materials (Cu) to increase plasticity [7, 8]. For example, 950 nm-diameter micropillars constituted of ZrCu/Cu amorphous/crystalline 10 nm-thick sublayers reports high $\sigma_y$ typical of Zr-based MGs (~2.4 GPa), while deforming homogeneously above 10% without SB propagation [7].

A second strategy to increase the size effect range on plasticity in MGs involves the control of their atomic and nanostructure through the introduction of local heterogeneities, altering the local chemistry/order, delaying the catastrophic failure due to SB propagation which are, therefore, activated at larger intrinsic sizes as a result of enhanced defect probability in large volumes [9]. This is the case of cluster-assembled TFMGs [10] and nanoglasses particles with reduced coherent domain size as observed in X-ray and synchrotron X-ray scattering. This structure with large number of nanointerfaces, local chemical fluctuations and high free volume content significantly affect the mechanical behavior, managing to accommodate local shear events and hindering the propagation of macroscopic SBs, promoting plasticity while keeping a high yield strength as a result for the disordered atomic structure. As an example, compression test of 1 µm-diameter micropillars of Zr<sub>50</sub>Cu<sub>50</sub> nanoglasses show up to 30% deformation, while maintaining a $\sigma_y$ >2.4 GPa and preventing the propagation of the catastrophic SBs in combination with homogeneous deformation [11].

Within this context, Pulsed Laser Deposition (PLD) stands among physical vapor deposition techniques for the possibility to synthesize novel phases, exploiting its extremely out-of-equilibrium features [12] in combination with the ease of fabrication of unique microstructures (i.e. compact/nanogranular [10] or ultrafine nanolaminates [13]) with tailored local heterogeneities by simply adjusting the deposition parameters combined with fast deposition rate for a rapid production of coatings. While PLD has mostly been used for oxides or ceramic materials [14, 15], recently, it has been implemented for the synthesis of novel metallic film with enhanced and tailored mechanical performance. As an example, nanogranular ZrCu TFMGs (thickness ~0.15 µm, cross-section ~0.06 µm²) report high $\sigma_y$ (~3 GPa) and homogeneous deformation > 9% during in situ TEM tensile tests, while avoiding the propagation of SBs [10].
Nevertheless, the synthesis of TFMGs by PLD with tailored composition and nanostructure is still at its infancy, and dedicated experiments are needed to understand the relationship with mechanical response, especially focusing on the mechanical behavior of larger size samples (thickness ≥ µm scale, cross-section > few µm²). This is a crucial step toward developing advanced TFMGs with tunable and optimal balance of strength/ductility, with extended mechanical size effects range to the micrometer scale and beyond, making them highly suitable for applications such as structural coatings. Specifically, the activation of mechanical size size effects at larger dimensions would be particularly beneficial for wear-resistant applications as a thicker protection layer would last longer extending the substrate material's service live.

In this work, we use PLD to synthesize TFMGs with tailored nanostructure (compact and nanogranular) and large thicknesses of 800 nm and 3 µm, focusing on model glass formers, namely Zr₅₀Cu₅₀ and Zr₄₆Cu₄₆Al₈, with Al selected for its ability to increase the strength and plasticity of Zr₅₀Cu₅₀ MGs [16, 17]. We then investigate the entire mechanical behavior, including compression test of micropillars tested at different strain rates from 0.001 up to 10 s⁻¹ and tensile properties on polymeric substrate, encompassing a large panel of structural and micromechanical characterization techniques. This enables us to correlate the local chemistry and nanostructure with the entire mechanical behavior, identifying the atomic and structural elements which enable them to extend the mechanical size effect range, leading to a large yield strength/plasticity balance up to few micrometers thickness (cross-section > 1.32 µm²).

Synchrotron X-ray scattering shows that PLD-deposited TFMGs have a unique and ultra-dense atomic structure with shorter interatomic bonds with respect to their sputtered counterparts. Moreover, nanogranular TFMGs report shorter coherence domains due to their cluster-assembled growth (clusters' dimensions ~0.7-1 nm) and a heterogeneous structure with compositional Zr-rich and Cu-rich fluctuations and high density of nanointerfaces. This results in unique mechanical behavior, showing that compression tests of 1.3 x 3 µm (diameter x thickness) micropillars can achieve large σ<sub>y</sub>>3 GPa with the propagation of SBs propagation occurring at strains >3%, for compact TFMGs. Such balance is even exceeded for nanogranular TFMGs maintaining a high σ<sub>y</sub> (>3 GPa), together with a homogeneous deformation >6% prior to shear banding.
We show that our PLD-TFMGs achieve significantly higher strength when compared with other MGs of the same dimensions, achieving values comparable to those observed in strongly confined MG specimens (pillar diameter < 550 nm, cross-section < 0.25 µm²). Finally, we show that such properties are kept during high-speed compression tests, while a high cracking resistance (>2% strain) is observed during tensile tests on polymeric substrate.

Overall, our results will pave the way for the use of PLD as a new synthesis route to fabricate TFMGs with tailored local heterogeneities and nanostructure, reporting a panel of large and tunable mechanical properties and high performances which are maintained for larger specimens' size, pushing the boundaries of mechanical size effects threshold with clear interest for applications such as protective coatings, flexible devices and microelectronics.

## 2. Materials and methods

### 2.1 Deposition of TFMGs by Pulsed Laser Deposition (PLD)

Zr₅₀Cu₅₀ and Zr₄₆Cu₄₆Al₈ (% at.) TFMGs were deposited by ablating an as-cast Zr₅₀Cu₅₀ (99.99% pure, produced by Mateck®) and a Zr₄₆Cu₄₆Al₈ target (99.95% pure, produced by Kurt J. Lesker®) with a ns-pulsed laser (Nd:YAG, 1st harmonic, λ = 1064 nm, repetition rate 10 Hz, pulse duration 5-7 ns), pulse energy equal to 1650 mJ, focused on the target to obtain a fluence of 14 J/cm². Si (100) and polyimide Kapton® were used as substrates after being cleaned in an ultrasonic bath in isopropanol. The substrates were mounted on a rotating substrate holder with a fixed distance from the target equal to 70 mm. Prior to the deposition the chamber was pumped down at 1x10⁻³ Pa vacuum level, then all depositions were performed at *T*room or in a pure He atmosphere at 5 Pa in order to deposit compact and cluster-assembled nanogranular TFMGs respectively, as described in Ref. [13]. 800 nm thick samples were used for structural characterizations by SEM, XRD, XRR and synchrotron X-ray total scattering measurements and mechanical characterizations by BLS, nanoindentation and tensile tests on polymer substrates while 3 µm thick samples were used for micropillar compression.
## 2.2 Structural characterization and ab initio molecular dynamics

A field emission scanning electron microscope (FE-SEM) Zeiss Supra 40 equipped with a Bruker energy dispersive X-ray spectroscopy (EDX) was used to perform morphological and elemental characterization. X-ray Diffraction (XRD) has been carried out with an Inel EQUINOX 1000 diffractometer equipped with a curved detector and Cu-Kα1 radiation (λ = 1.5406 Å) and the angle of incidence of the X-ray beam was 2°. Average interatomic *d*<sub>av</sub> could be extracted from the hump peak position using Erhenfest equation [18]. X-ray reflectivity (XRR) scans were obtained with a Seifert XRD 3000 diffractometer operating at 40 kV and 40 mA using a Cu Kα source. The scans were carried out between 0.2° and 2.2° with a step of 0.05° to determine the mass density of the films.

The pair distribution functions of the TFMGs were obtained from X-ray total scattering measurements performing dedicated experiments at the French radiation synchrotron facility (SOLEIL) on the Diffabs beamline in reflection mode. The samples consisted in thin film deposited at the surface of Si (100) (1 x 1 cm²). For all experiments, the monochromatic beam's energy was fixed at 18 keV, corresponding to a wavelength of about 0.68 Å. The incidence angle was fixed to 2°. A CirPAD two-dimensional detector [19] was mounted at about 64 cm from the sample surface to acquire 2D diffractograms in reflection mode which were then converted to 1D diffractogram using a home-made software developed at Diffabs Beamline. The diffractograms were obtained with 2θ range from 3.75 up to 139.40° which corresponds to a Q scattering vector range of 0.6-17.4 Å⁻¹. The total scattering pattern of the Si (100) substrate was acquired in the same conditions for background subtraction. Data reduction and correction were performed using the PDFgetx3 software [20] to obtain the experimental reduced PDF *G*<sub>exp</sub>(r). During the correction procedure, the background subtraction and Fourier transformed Q-range were finely adjusted to maximize the TFMG signal. The experimental PDF *g*<sub>exp</sub>(r) was obtained from *G*<sub>exp</sub>(r) values for comparison with Ab Initio Molecular Dynamics (AIMD) calculations.

AIMD were performed to simulate the amorphous structure and compared with the experimental structural data (mass density, average interatomic distance and PDF). First-principles calculations were performed with the Vienna Ab initio Simulation Package (VASP) [21]. 256 atoms of ZrCuAl<sub>x</sub> (x = 0, 8 % at.)
were placed in a cubic box with periodic boundary conditions. The two alloys were first melted at 2500 K and
maintained at this temperature for at least 12 ps using the isobaric-isothermal (NPT) ensemble to obtain a
homogeneous molten alloy. The system was then continuously quenched from 2500 down to 300 K at a
cooling rate of $733 \times 10^{13}$ K/s (0.3 $\times 10^{-15}$ s time step and 1000 steps) using the isovolumic-isothermal (NVT)
ensemble and then relaxed at 300 K using NPT for at least 11 ps. Complementary information regarding the
synchrotron characterization can be found in Section S1 of the supporting information (SI).

## 2.3 Mechanical characterization

Brillouin light scattering (BLS) has been carried out in air using a green laser source ($\lambda$ = 532 nm) to obtain the elastic properties of the films including the shear (G) and elastic (E) moduli and the Poisson's ratio (v), using the same procedure for isotropic materials described in Ref. [22]. Examples of the raw BLS spectra and fits can be found in Section S2 of the SI.

The mechanical properties and the plastic behavior have been extracted by nanoindentation and *in situ* SEM micropillar compression. Nanoindentation has been performed with a KLA G200 with a Berkovich diamond indenter operated in continuous stiffness measurement (CSM). Prior to the measurements, the area function was calibrated, performing nanoindentation on standard fused silica. For each sample, 25 indents were performed, and the Oliver and Pharr model was applied to extract *H* at indentation depth equal to 10 % of the film thickness [23, 24]. Examples of the nanoindentation curves can be found in Section S3 of the SI.

*In situ* SEM micropillar compression tests were performed with an Alemnis ASA nanoindenter with an 1.3 $\times$ 3 µm (diameter $\times$ thickness) micropillars compressed with a single step of deformation reaching 25% total deformation, using a 5 µm diameter flat punch tip. For each strain rate 5 pillars were tested in the SEM while the electron beam was in continuous scanning mode (5 kV). Examples of the stress-strain curves presented in Section S4 of the SI. The viscoplastic behavior in compression was examined through compression at different $\epsilon$ from $10^{-3}$ (quasi-static loading) up to $10^{-1}$ s$^{-1}$. The values of the stress drop caused by SB propagation have been measured by a close examination of the stress-strain curves after yielding (the raw results can be found in section S5 of the SI).
Tensile tests were performed on films deposited on polyimide Kapton® substrates, using a 300 N Deben tensile machine combined with a Keyence confocal microscope. The Kapton® substrates were 4 cm long and 4 mm wide and the films were deposited in the center of the substrate in an area 1 cm long and 4 mm wide to avoid artifacts of the measurement related to inhomogeneous thickness of the films. The tensile speed was set equal to 0.5 mm/min and the sample length was 40 mm corresponding to a strain rate of ~2.1x10⁻⁴ s⁻¹. A 10x objective was used for visualizing the evolution of the crack density as a function of the strain. The linear crack density, which measures the number of cracks over the unit of length perpendicular to their propagation, was calculated by measuring the number of cracks (orthogonal to the loading direction) over the area of the sample (1412 µm x 1059 µm). From these tests it is possible to obtain the crack onset strain (COS), when the first crack appears. Details regarding the model used to extrapolate the residual stresses from the experimental data can be found in Section S6 of the SI.

## 3. Results

### 3.1 Structural properties

Fig. 1a shows the SEM cross-sections after mechanical cleavage of the Zr₅₀Cu₅₀ and Zr₄₆Cu₄₆Al₈ TFMGs deposited in vacuum and in 5 Pa He. All samples present the corrugations typical of MGs and related to the propagation of shear bands (SBs) during fracture [25]. However, the width of the corrugations on the surface of the cross-sections, decrease with the addition of Al from ~70 ± 5 down to ~60 ± 5 nm for, respectively, compact Zr₅₀Cu₅₀ and Zr₄₆Cu₄₆Al₈ in line with literature values for Zr-based TFMGs [5, 16, 26]. A reduction of the size of the corrugations can be related to a change in the cavitation mechanism within the amorphous structure during cleavage due to the presence of Al favoring a less localized nucleation of SBs [17]. The transition from atom-by-atom (compact film, schematic of Fig. 1b), to cluster-assembled growth, (nanogranular films schematic of Fig. 1d), occurs when the deposition pressure exceed 1-5 Pa [10], reducing the size of the corrugations down to ~60 ± 4 and 40 ± 3 nm for nanogranular Zr₅₀Cu₅₀ and Zr₄₆Cu₄₆Al₈, respectively. This can be connected to the small local clustering, favoring the delocalization of SBs during the
cleavage similar to nanoglasses [11, 27]. The formation of local clustering is in line with our previous works and has been observed more in detail with HRTEM characterization in Refs. [10, 13].

<figure><img src="image_3.png" /><figcaption>**Figure 1:** (a, c) SEM cross-section micrographs of, respectively, compact and nanogranular Zr<sub>50</sub>Cu<sub>50</sub> and Zr<sub>46</sub>Cu<sub>46</sub>Al<sub>8</sub> TFMGs deposited by Pulsed Laser Deposition (PLD). (b, d) schematics depicting, respectively, atom-by-atom growth (formation of compact films) and cluster-assembled growth (formation of nanogranular films).</figcaption></figure>

The XRD patterns of the compact and nanogranular Zr<sub>50</sub>Cu<sub>50</sub> and Zr<sub>46</sub>Cu<sub>46</sub>Al<sub>8</sub> films are presented in Fig. 2a. All TFMGs present a single broad hump centered at ~37.7°, corresponding to an interatomic distance of 2.93 Å (Ehrenfest equation), similar for bulk and sputter-deposited ZrCu and ZrCuAl [13, 16]. The addition of a small quantity of Al (up to 12 %at.) doesn't affect the XRD patterns as observed for sputter-deposited counterparts [16]. However, the transition from compact to nanogranular nanostructure increases the full width half maximum (FWHM) of both Zr<sub>50</sub>Cu<sub>50</sub> and Zr<sub>46</sub>Cu<sub>46</sub>Al<sub>8</sub> TFMGs from, respectively, ~7.5 up to 11° due to the
smaller coherence domain of the structure with respect to compact TFMGs due to the high amount of disorder connected to the cluster-assembled structure of nanogranular TFMGs, see Refs. [10, 13].

The atomic local structure has been then probed using atomic pair distribution function (PDF) method to track the structural changes due to the different compositions and nanostructures. Fig. 2b shows the experimental PDFs ($g_{exp}(r)$) of the compact and nanogranular $Zr_{50}Cu_{50}$ and $Zr_{46}Cu_{46}Al_{8}$ TFMGs. All the experimental PDF show the same general pattern with a main contribution ~2.7 - 2.8 Å together with few less intense contributions visible until 10-12 Å typical of MGs and amorphous structure without long range order and whose trends are confirmed by the S(Q) patterns present in Fig. S1.

Fig. 2c (and Figure S2) shows that the experimental and MD calculated $g_{exp}(r)$ for compact $Zr_{50}Cu_{50}$ and $Zr_{46}Cu_{46}Al_{8}$ TFMGs are in very good agreement, indicating that the local structure of the compact TFMG is similar to its bulk counterpart, enabling a rigorous discussion about the evolution of interatomic distances in TFMGs. Compact $Zr_{50}Cu_{50}$ has average interatomic distance of 2.79 Å which slightly decreases down to 2.75 Å with the addition of Al due to the formation of shorter Zr-Al and Cu-Al bonds in agreement to literature [16], see also Fig. S3 of the SI. On the other hand, the experimental PDF of the nanogranular $Zr_{50}Cu_{50}$ sample shows a peak with interatomic distance equal to 2.70 Å (which goes down to 2.68 Å for nanogranular $Zr_{46}Cu_{46}Al_{8}$) in agreement with the average Zr-Cu bonds and a shoulder at 3.30 Å which is related to the Zr-Zr bonds (dashed line in Fig. 2b) which is also present in compact $Zr_{50}Cu_{50}$ with a lower intensity due to the limited atomic scale chemical enrichments related to atom-by-atom growth. The intensity of the contribution at 3.30 Å is greatly reduced for nanogranular $Zr_{46}Cu_{46}Al_{8}$ as a result of the formation Zr-Al bonds in favor of Zr-Zr bonds. The shift towards smaller interatomic distances for nanogranular TFMGs can be correlated with the formation of chemical heterogeneities between the core and the shell of the clusters which could be associated with a slightly Cu-rich core and Zr-rich shell [22, 28]. The interatomic distance of PLD-deposited TFMGs are slightly smaller than the interatomic distances observed in Zr-based bulk MGs [29] and in the sputtered ZrCu counterparts (2.82 Å, see Figure S3), indicating a more densely packed structure as expected from the high deposition energies associated with PLD (10-15 eV vs 1-5 eV for sputtering) [30, 31].
Fig. 2d highlights the differences of coherence domains size between compact and nanogranular nanostructures by focusing on Zr<sub>50</sub>Cu<sub>50</sub> PDFs extracted from Fig. 2b. Compact Zr<sub>50</sub>Cu<sub>50</sub> has a large coherence domain showing 3 shells of interaction centered ~2.80, 4.91 and 7.50 while its nanogranular counterparts possess only two shells of interaction centered ~2.70 and 4.98 Å. Such reduction is related to clustering formation during growth with the core confined in the 6.4 – 6.6 Å (dashed line in Fig. 2d), indicating a large amount of nanointerfaces and free volume with respect to compact counterparts.

<figure><img src="image_4.png" /><figcaption>**Figure 2.** (a) XRD diffractograms and (b) pair distribution functions (PDFs) of compact and nanogranular Zr<sub>50</sub>Cu<sub>50</sub> and Zr<sub>46</sub>Cu<sub>46</sub>Al<sub>8</sub> TFMGs. (c) The PDF of compact Zr<sub>50</sub>Cu<sub>50</sub> and the result from AIMD simulation. (d) Detailed comparison between the PDFs of compact and nanogranular Zr<sub>50</sub>Cu<sub>50</sub> TFMGs, highlighting the different number of interaction shells.</figcaption></figure>

Table 1 shows the mass density ($ρ$) of TFMGs deposited by PLD with different Al content and nanostructure.

The addition of Al results in a decrement of $ρ$ from 8350 down to 7920 $Kg/m^3$ for compact Zr<sub>50</sub>Cu<sub>50</sub> and
Zr₄₆Cu₄₆Al₈ as a result of the Al incorporation (ρ = 2700 Kg/m³), substituting heavier Zr (ρ = 6505 Kg/m³) and Cu (ρ = 8940 Kg/m³) atoms. Furthermore, the compact TFMGs have a ρ ~ 10% higher than MS counterparts, which is ~7400 and 7170 Kg/m³ for Zr₅₀Cu₅₀ and Zr₄₆Cu₄₆Al₈, respectively [16]. This is related to the high energies involved in PLD process, which increase the kinetic energy and momentum of the ablated species, resulting in a more closely packed atomic structure [30], as already observed for carbon [32] and oxide [14] films. This is also in agreement with the previous synchrotron study reporting a reduction in the average atomic distances.

The transition from compact to nanogranular nanostructure reduces the ρ down to 7980 and 7100 Kg/m³ for nanogranular Zr₅₀Cu₅₀ and Zr₄₆Cu₄₆Al₈, respectively. This is connected to the cluster-assembled growth characterized by high density of nanointerfaces and larger free volume content similarly to nanoglasses [10]-[13]. Moreover, the transition from compact to nanogranular nanostructure does not affect linearly the ρ trend with the nanogranular Zr₅₀Cu₅₀ being ~4.5% less dense than its compact counterpart, while nanogranular Zr₄₆Cu₄₆Al₈ is ~11% less dense than compact Zr₄₆Cu₄₆Al₈. The presence of Al influences the cluster growth by reducing the amount of Zr-Zr bonds in favor of Zr-Al bonds, in agreement with the disappearance of the Zr-Zr shoulder (Fig. 2b) and the increment of free volume at the interface shells of the nanogranular structure.

<table><caption>**Table 1:** Mass density of compact and nanogranular Zr₅₀Cu₅₀ and Zr₄₆Cu₄₆Al₈ TFMGs.</caption><thead><tr><td>Sample</td><td>Mass Density (Kg/m³)</td></tr></thead><tbody><tr><td>Compact Zr<sub>50</sub>Cu<sub>50</sub></td><td>8350 ± 50</td></tr><tr><td>Nanogranular Zr<sub>50</sub>Cu<sub>50</sub></td><td>7980 ± 50</td></tr><tr><td>Compact Zr<sub>46</sub>Cu<sub>46</sub>Al<sub>8</sub></td><td>7920 ± 50</td></tr><tr><td>Nanogranular Zr<sub>46</sub>Cu<sub>46</sub>Al<sub>8</sub></td><td>7100 ± 50</td></tr></tbody></table>

## 3.2 Elastic properties and hardness

Table 2 shows the elastic modulus (E) and hardness (H) of compact and nanogranular Zr₅₀Cu₅₀ and Zr₄₆Cu₄₆Al₈ TFMGs which have been investigated with Brillouin Light Scattering and nanoindentation. Firstly, it can be observed that the differences between the E and H of compact and nanogranular TFMGs with the same
composition are very low, suggesting that the change in nanostructure does not have a major effect on the mechanical properties, which instead are mainly governed by the strength of the chemical bonds. The *E* and *H* values of the PLD-deposited TFMGs, respectively 121 and 8.6 GPa and 109 and 7.8 GPa for Zr<sub>50</sub>Cu<sub>50</sub> and Zr<sub>46</sub>Cu<sub>46</sub>Al<sub>8</sub>, are ~30% larger than the MS counterparts, which are equal to 85 and 6.7 GPa and 88 and 7.3 GPa for Zr<sub>50</sub>Cu<sub>50</sub> to Zr<sub>46</sub>Cu<sub>46</sub>Al<sub>8</sub>, respectively [29, 33] as a result of higher mass density and high kinetic energies involved in the PLD process [13, 34]. Finally, the value of *E* obtained from the fit of the loading curve of Zr<sub>50</sub>Cu<sub>50</sub> and Zr<sub>46</sub>Cu<sub>46</sub>Al<sub>8</sub> TFMGs during micropillar compression (Section 3.3) are, respectively, equal to ~108 ± 10 and 90 ± 15 GPa for compact TFMGs and 109 ± 11 and 92 ± 13 GPa for nanogranular TFMGs. Such values are close to the ones obtained by optoacoustic techniques (Table 2) and within the standard deviation expected from the measurement by micropillar compression tests.

<table><caption>**Table 2:** Elastic modulus and hardness of compact and nanogranular Zr<sub>50</sub>Cu<sub>50</sub> and Zr<sub>46</sub>Cu<sub>46</sub>Al<sub>8</sub> TFMGs.</caption><thead><tr><td>Sample</td><td>Elastic modulus (BLS) (GPa)</td><td>Hardness (nanoindentation) (GPa)</td></tr></thead><tbody><tr><td>Compact Zr<sub>50</sub>Cu<sub>50</sub></td><td>121 ± 2</td><td>8.6 ± 0.2</td></tr><tr><td>Nanogranular Zr<sub>50</sub>Cu<sub>50</sub></td><td>121.5 ± 2</td><td>8.3 ± 0.2</td></tr><tr><td>Compact Zr<sub>46</sub>Cu<sub>46</sub>Al<sub>8</sub></td><td>110 ± 2</td><td>7.7 ± 0.2</td></tr><tr><td>Nanogranular Zr<sub>46</sub>Cu<sub>46</sub>Al<sub>8</sub></td><td>112 ± 2</td><td>7.9 ± 0.2</td></tr></tbody></table>

### 3.3 Plastic behavior by micropillar compression

The strength and plasticity of the PLD TFMGs has been investigated by compression tests of micropillars in order to explore the mechanical response under uniaxial compression at a standard loading rate (ε̇) equal to 10⁻² s⁻¹. Fig. 3a shows the engineering stress-strain compression curves of compact Zr<sub>50</sub>Cu<sub>50</sub> and Zr<sub>46</sub>Cu<sub>46</sub>Al<sub>8</sub> micropillars. Compact Zr<sub>50</sub>Cu<sub>50</sub> TFMGs deform elastically up to 3% engineering strain upon which the yield strength (σ<sub>y</sub>) is reached (3.2 ± 0.1 GPa). This is followed by homogenous plastic deformation up to 3.2 ± 0.2 %, when few large SBs propagate, Fig. 3b. On the other hand, compact Zr<sub>46</sub>Cu<sub>46</sub>Al<sub>8</sub> TFMGs deform elastically up to 3.7%, reaching yielding at 2.8 ± 0.1 GPa, followed by homogenous deformation up to 4.1 ± 0.2 % upon which SB failure occurs. This trend agrees with the hardness reported in Table 2, while the improved plasticity in Zr<sub>46</sub>Cu<sub>46</sub>Al<sub>8</sub> is related to the larger free volume and formation of strong Zr-Al bonds, leading to the formation
of multiple, but non percolative, SBs [17, 35]. Lastly, it can be observed that the $σ_y$ and the plastic strain necessary for the formation of SBs in PLD-deposited TFMGs are larger than in MS counterparts, reaching the maximum of $σ_y$ around 2.5 GPa and plasticity up to ~2% - 3% strain [36, 37]. This is due to the denser atomic structure and larger mechanical properties of the PLD TFMGs, see Refs. [10, 13].

Figure 4a shows the engineering stress-strain curves of the nanogranular $Zr_{50}Cu_{50}$ and $Zr_{46}Cu_{46}Al_8$ micropillars. At 3% elastic deformation, nanogranular $Zr_{50}Cu_{50}$ reach yielding ($σ_y$ = 2.9 ± 0.1 GPa), slightly below the compact counterpart and in line with the *H* values (Table 2). This is connected to the higher amount of free volume, enabling plastic deformation at lower stresses [38, 39]. After yielding, the deformation continues homogeneously up to 6.0 ± 0.2% when the first SB develops (Fig. 4b). Nanogranular ZrCu TFMGs have a lower number and more superficial SBs with respect to their compact counterpart (Fig. 3b) similarly to what observed in cluster-assembled ZrCu deformed in tension [10] and ultrafine ZrCu/Al nanolaminates [13].

<figure><img src="image_5.png" /><figcaption>Figure 3: (a) Engineering stress-strain curves from micropillar compression of compact $Zr_{50}Cu_{50}$ and $Zr_{46}Cu_{46}Al_8$ TFMGs deposited by PLD and (b, c) the corresponding post-mortem SEM images.</figcaption></figure>

On the other hand, nanogranular $Zr_{46}Cu_{46}Al_8$ micropillars deform elastically up to 4% with $σ_y$ = 3.2 ± 0.1 GPa, followed by homogeneous deformation up to 6.7 ± 0.2%, when non-percolative SBs nucleate on the surface of the micropillar, Fig. 4c. The increase of $σ_y$ vs compact counterparts, can be related to the presence of Al based chemical heterogeneities with stronger atomic bonds between the core and the boundaries of the clusters, Fig. 2d. This leads to an increment of the stress necessary to reach yield and induce SB
propagation. This phenomenon similar to what has been observed for Fe₂₅Sc₇₅ nanoglasses, reporting local
chemical heterogeneities with stronger atomic bonds increasing σ_y vs monolithic counterpart [40].

<figure><img src="image_6.png" /><figcaption>**Figure 4:** (a) Engineering stress-strain curves from micropillar compression of nanogranular Zr₅₀Cu₅₀ and Zr₄₆Cu₄₆Al₈ TFMGs and (b, c) the corresponding post-mortem SEM images.</figcaption></figure>

**3.4 Viscoplastic behavior by micropillar compression**

The mechanical behavior of compact and nanogranular Zr₄₆Cu₄₆Al₈ TFMGs, presenting a larger plasticity due
to the addition of Al, have been further studied by micropillar compression at different strain rates to
investigate the viscoplastic behavior and observe if the mechanical properties are kept at high loading rates.
Compact Zr₄₆Cu₄₆Al₈ (Fig. 5a) TFMGs show a decrement of σ_y increasing ε̇, from 2.9 GPa (ε̇ = 10⁻³ s⁻¹) down to
2.55 GPa (ε̇ = 10 s⁻¹), resulting in a negative strain rate sensitivity coefficient (m) equal to -0.014 ± 0.0015, as
shown in Fig. 5c. Negative m values have already been observed for different MGs and are associated with
serrated plastic flow driven by the relaxation of shear transformation zones (STZs) and subsequential
propagation of SBs. During this process, mature SBs propagate, aided by small inhomogeneities within the
amorphous structure which become active once a critical strain is reached releasing the energy accumulated
from compression. As the ε̇ increases, there is less time to accumulate elastic energy before reaching the
critical strain, leading to a reduction of the yield strength [41, 42].

Fig. 5d shows the stress drop caused by the propagation of SBs during compression test calculated by
averaging the values of the single drops per curve, more details in Figure S8. The stress drop decreases with
higher ε̇ decreasing the amount of elastic energy necessary for the propagation of SBs [43, 44]. A similar
phenomenon is observed in the works of Schuh *et al.* [45] and Brognara *et al.* for [22], respectively, during nanoindentation of ZrCu bulk and TFMGs in which the magnitude of the strain bursts (that are directly proportional to the values of the stress drops) decrease at higher loading rate. The value of these stress drops from ~250 MPa (quasi-static loading) down to ~100 MPa for $\dot{\epsilon}$ equal to 0.1 s⁻¹, similar to literature values for Zr-based MGs [41]. Finally, the absence of stress drops for $\dot{\epsilon} > 0.1$ s⁻¹ is caused by the higher floor noise.

<figure><img src="image_7.png" /><figcaption>**Figure 5:** Engineering stress-strain curves from micropillar compression interrupted after yielding of (a) compact and (b) nanogranular Zr₄₆Cu₄₆Al₈ TFMGs with different $\dot{\epsilon}$ and associated post-mortem SEM images. The scale bar is equal to 1 µm. (c, d) Evolution of $\sigma_y$ and stress drop during shear banding vs $\dot{\epsilon}$. The stress drops values for $\dot{\epsilon} > 0.1$ s⁻¹ are absent due to the increased noise measurement at high-speed.</figcaption></figure>
Nanogranular Zr₄₆Cu₄₆Al₈ TFMGs have an opposite behavior with respect to their compact counterpart with σ<sub>y</sub> increases with ἐ from 3.08 GPa ( = 10⁻³ s⁻¹) up to 3.50 GPa (em = 10 s⁻¹), resulting in a positive *m* = 0.014 ± 0.0015 (Fig. 5 b,c). The transition from a negative to a positive *m* is associated with a change in the deformation mechanisms and propagation of SBs. Specifically, positive *m* values are connected to the activation of multiple and smaller proto SBs whose mutual interaction result in an increment of the energy necessary for propagation at high ἐ [46]. This is also supported by the fact that the stress drop values of nanogranular Zr₄₆Cu₄₆Al₈ TFMGs are much lower (~50 MPa) than their compact counterparts (~250 MPa) and are not influenced by the change of ἐ, Fig. 5d and Refs. [41, 46]. Here, the propagation of mature SBs does not happen and instead the deformation is mediated by the propagation of embryonic proto-SBs activated at the interfaces between different clusters. The *post-mortem* images in Fig. 5b show that the micropillars have some cracks which superficially resemble to SBs, similarly to what have been observed on *post-mortem* TEM characterization of nanogranular ZrCu/Al ultrafine nanolaminates deformed in compression in which no percolated cracks are observed, reporting an homogeneous deformation without marked stress drops [13].

When compared with literature the values of *m* of PLD films appear to be slightly smaller than their sputtered counterparts (0.025) observed in Ref. [16] which can be connected to a difference in the amount of free volume between the films produced by different deposition methods. In the article of Bhattacharyya *et al.* [47] it is shown that by decreasing the free volume of a metallic glass sample the value of *m* increases while *m* decreases if the free volume is increased. It follows that sputtered films have less free volume than PLD films, especially nanogranular ones.

**3.5 Fracture behavior**

Uniaxial tensile tests of TFMGs deposited on Kapton® substrates were performed to investigate their
mechanical response and fracture in tension when coupled with a flexible substrate. Figure 6 reports the
evolution of the crack density as a function of the engineering strain of the TFMGs with different composition
and nanostructure. Compact TFMGs have a the crack onset strain (COS) which increase with the addition of
Al from 1.30 ± 0.05 up to 2.03 ± 0.05 % for, respectively, compact Zr<sub>50</sub>Cu<sub>50</sub> and Zr<sub>46</sub>Cu<sub>46</sub>Al<sub>8</sub> (Figure 6a). These
values are higher than the corresponding MS films with the same thickness (800 nm, 1% for Zr₅₀Cu₅₀ and 1.7% for Zr₄₆Cu₄₆Al₈), Ref. [16]. This can be related to the combination of the high strength of PLD TFMGs and compressive residual stresses (-270 ± 50 MPa) for compact PLD TFMGs (see SI, Section 6) vs sputtered counterparts (-60 ± 20 MPa) [48], delaying crack onset formation [22]. On the other hand, nanogranular TFMGs show a decrement of the COS from 1.04 (Zr₅₀Cu₅₀) down to 0.57 % (Zr₄₆Cu₄₆Al₈), accompanied by the formation of straight cracks without deviations and blocking phenomena (Figure 6b). The reduced crack resistance of the nanogranular TFMGs can be related to the cluster-assembled structure, with the presence of atomic interfaces [10, 11, 13], since both compact and nanogranular TFMGs have similar compressive stresses, respectively -270 ± 50 and -290 ± 55 MPa, with no variation with the chemical composition (the exact values are reported in Table T2 of SI). In literature, the presence of local interfaces favor the crack initiation such as for nanocolumnar Zr₅₂Cu₄₈ reported in Ref. [22] (COS = 0.5%, lower than the homogeneous counterparts, 1%) resulting in a more fragile behavior. The local density and chemistry fluctuations of nanogranular TFMGs can have a similar effect and favor the fragmentation during tensile loading.

<figure><img src="image_8.png" /><figcaption>**Figure 6:** (a) Crack density vs engineering strain curves of compact and nanogranular Zr₅₀Cu₅₀ and Zr₄₆Cu₄₆Al₈ TFMGs during tensile test and (b) post-mortem images. The yellow arrow indicates the loading direction.</figcaption></figure>
<table><caption>**Table 3:** Crack onset strain (COS) of compact and nanogranular Zr<sub>50</sub>Cu<sub>50</sub> and Zr<sub>46</sub>Cu<sub>46</sub>Al<sub>8</sub> TFMGs obtained from tensile tests Kapton® substrates.</caption><thead><tr><td>Sample</td><td>Crack onset strain (%)</td></tr></thead><tbody><tr><td>Compact Zr<sub>50</sub>Cu<sub>50</sub></td><td>1.3±0.05</td></tr><tr><td>Nanogranular Zr<sub>50</sub>Cu<sub>50</sub></td><td>1.04±0.05</td></tr><tr><td>Compact Zr<sub>46</sub>Cu<sub>46</sub>Al<sub>8</sub></td><td>2.03±0.05</td></tr><tr><td>Nanogranular Zr<sub>46</sub>Cu<sub>46</sub>Al<sub>8</sub></td><td>0.57±0.05</td></tr></tbody></table>

## 4. Discussion

In this work, we synthesized novel nanostructured TFMGs with tailored local heterogeneities, exploiting the large potential of pulsed laser deposition (PLD), with the objective of extending the mechanical size effect range ensuring a large combination of mutually exclusive mechanical properties (yield strength and ductility) for intrinsically large dimensions.

The investigation of the atomic structure by synchrotron total scattering and PDF reveals that the addition of Al results in a formation of shorter Zr-Al bonds [16] for both compact and nanogranular TFMGs. Nanogranular TFMGs show PDFs with two peaks of interatomic distances associated with local chemical heterogeneities and a smaller coherence domain vs compact counterpart (Fig. 2d). This is the signature of the cluster-assembled structure with high free volume and nanointerfaces [11]. The higher atomic packing of PLD-deposited TFMGs with respect to their sputtered MS counterparts found in literature results in higher $\rho$ and consequently higher $E$ and $H$, up to 121 and 8.6 GPa respectively, (25/30 % higher than sputtered counterparts [16][49]) and large $\sigma_y$ (> 2.8 GPa) in compression. However, the fracture strain is strongly dependent on the nanostructure. Specifically, compact PLD TFMGs can reach a fracture strain up to 4.1% for Zr<sub>46</sub>Cu<sub>46</sub>Al<sub>8</sub> due to the presence of short Zr-Al bonds, decreasing the atomic mobility and hindering the propagation of SBs [35]. On the other hand, nanogranular TFMGs fail at larger strains (~6.7%) due to their cluster-assembled structure, preventing the formation of mature SBs and promoting a more homogeneous deformation mediated by proto-SBs confined between the clusters' interfaces [50, 51]. Micropillars compression test at high $\dot{\epsilon}$ reveal that compact TFMGs have negative strain rate sensitivity ($m$), connected to the propagation of mature SBs similar to bulk and sputter-deposited MGs [42, 52]. On the contrary, the
positive *m* is for nanogranular counterparts, indicates that the deformation is mediated by multiple and smaller proto SBs, which, however, avoid the catastrophic failure [46].

To understand where PLD-deposited TFMGs fit within the literature and the extension of the mechanical size effects for model ZrCu glass formers, we compared the data of yield strength/plastic deformation as a function of the intrinsic size with other MGs and nanoglasses found in literature tested either in compression [11, 53-58] or in tension [6, 54], as shown in Fig. 7. Thick MGs fail in a brittle way ~3% deformation with the propagation of SBs (green area, Fig. 7). However, reducing the intrinsic size below a critical threshold (~500 - 550 nm) mechanical size effects are activated, hindering the formation of SBs and increasing the σ<sub>y</sub> towards the ideal limit (atomic decohesion), together with the strain at failure which can reach 10% - 15% [4, 6, 59] (purple area, Fig. 7). The nanoglasses shown in Figure 7 (orange areas) are characterized by large plastic deformation (up to 30%) without SB propagation thanks to their nanogranular structure while their σ<sub>y</sub> stays in the range of the other MGs found in literature.

<figure><img src="image_9.png" /><figcaption>**Figure 7:** Strength versus sample size (or cross-sectional area) in sub-micrometer MGs samples deformed in tension or compression, highlighting mechanical size effects activation for critical dimension <550 nm (dashed lines are guide for eyes). PLD-deposited TFMGs manage to extend mechanical size effects range with yield strength close to the theoretical limit and plastic deformation > 6 % for 1.3 µm diameter (1.32 µm² cross-section) specimens.</figcaption></figure>
The PLD-deposited TFMGs studied in this work (red area, Fig. 7) with a 1.3x3 µm (height x diameter) and 1.32 µm² cross-section report high σ<sub>y</sub> (>2.8 GPa) and plastic deformation (>6%) balance which would have required stronger geometrical confinement to be achieved (average diameter < 0.55 µm, cross-sectional area 0.25 µm²) for conventional MGs. This is associated with the unique features of PLD-deposited films with denser and more tightly packed atomic structure with smaller and stronger interatomic bonds. Moreover, thick PLD TFMGs have a σ<sub>y</sub> which is ~60-70% of the theoretical strength (atomic decohesion) for ZrCu/ZrCuAl MGs [49], calculated as E/20 from Ref. [1], while conventional MGs need to be strongly confined (average diameter < 0.55 µm) to achieve similar results, Fig. 7. In addition, compact Zr<sub>50</sub>Cu<sub>50</sub>Al<sub>8</sub> can deform up to 4.1%, and > 6% for nanogranular counterparts more than doubling the value for conventional MGs with similar size [55, 57]. This is connected to nanoclustering and high free volume content capable of preventing the propagation of mature SBs and pushing toward higher deformations the catastrophic failure [10, 13]. Furthermore, the very low value of m (~±0.014) for PLD-deposited TFMGs, when compared with Zr-based MGs tested by micropillar compression [41] (~-0.055), indicates a stable mechanical strength maintaining a large σ<sub>y</sub> > 2.5 GPa and nearly constant fracture strain >4% even when subjected to high strain rates. Finally, tensile tests confirm the extended plasticity of PLD TFMGs even in complex loading configurations on polymeric substrate, reaching 2% COS for compact Zr<sub>46</sub>Cu<sub>46</sub>Al<sub>8</sub> which results of being 18% more crack resistant than their literature counterparts [16].

**5. Conclusions**

In this work, we used PLD to deposit model TFMGs Zr<sub>50</sub>Cu<sub>50</sub> and Zr<sub>46</sub>Cu<sub>46</sub>Al<sub>8</sub> with compact and nanogranular
nanostructure aimed to extend their mechanical size effect range over different loading configurations. We
conducted a detailed investigation of how variations in local chemistry and nanostructure correlate with
mechanical behavior. Specifically, we show that:

* The addition of Al to compact and nanogranular TFMGs results in a reduction of the average interatomic distances (formation of strong Zr-Al and Cu-Al bonds), which enhance atomic packing. On
the other hand, nanogranular TFMGs have smaller coherence domains due to their cluster-assembled
growth with small (6.4 – 6.6 Å) cores.

* PLD-deposited TFMGs display higher values of mass density, *E* and *H* vs sputtered counterparts because of the extremely energetic conditions of the PLD process promoting denser atomic packing. This also explains the large *COS* (2%) for compact Zr<sub>46</sub>Cu<sub>46</sub>Al<sub>8</sub> during tensile test on polymeric substrate

* Micropillar compression of compact TFMGs reveals that they have a σ~y~ > 2.8 GPa and that the addition of Al increases the plasticity from 3.2% to 4.1%, while the nanogranular TFMGs can achieve strain at failure up to 6.7%, while keeping a superior σ~y~ ~3 GPa due to their cluster-assembled structure preventing the formation of mature SBs. Moreover, the absolute low value of *m* (~±0.014) indicates that PLD TFMGs have stable (and large) mechanical strength even when subjected to high ε, while the difference in sign of *m* is related with a difference on the deformation behavior with the formation of mature and smaller SBs, respectively, for compact and nanogranular TFMGs.

* The comparison with current literature of MGs shows that PLD TFMGs can achieve high σ~y~ combined with large plastic strain significantly above what is usually obtained for traditional MG families with equivalent dimensions. Specifically, 1.3 x 3 µm (height x diameter) PLD-deposited TFMGs possess a yield strength and strain and failure balance similar to sub-micron scale TFMGs but possessing diameter < 0.55 µm and cross-sectional area < 0.25 µm².

In summary, our work demonstrates that PLD can produce nanostructured TFMGs that combine high
mechanical strength with sustained plastic deformation even in large samples exceeding the dimensions in
which mechanical size effects typically emerge in conventional MGs. These findings pave the way for the
future use of TFMGs as structural materials in protective coatings for fatigue- and wear-resistant applications,
as well as in microelectronics, flexible devices and energy-related technologies.

**Appendix: Supporting information**

"Additional details regarding the characterization techniques including examples of raw data".
**Author contribution**

**Francesco Bignoli**: investigation, validation, formal analysis, data curation, writing - original draft, visualization. **Alex Lemarchand**: investigation, validation, data curation, writing - review & editing. **Szilvia Kalácska**: investigation, data curation, writing - review & editing. **Dominique Thiaudiere**: investigation. **Philippe Djemia**: investigation, validation, writing - review & editing. **Damien Faurie**: resources, supervision, writing - review & editing, funding acquisition. **Andrea Li Bassi**: resources, supervision, writing - review & editing, funding acquisition. **Matteo Ghidelli**: conceptualization, resources, supervision, writing - review & editing, project management, funding acquisition.

**Declaration of Competing Interest**

The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

**Acknowledgements**

M. Ghidelli, A. Li Bassi and D. Faurie acknowledge the financial support of the Université Franco Italienne (UFI) for the funding of the PhD scholarship of F. Bignoli through a Vinci Cap. III Grant (C3-2286). M. Ghidelli acknowledges the financial support of the French National Research Agency (ANR) projects MICRO-HEAs (N° ANR-21-CE08-0003-01) and EGlass (N° ANR-22-CE92-0026-01). S. Kalácska was funded by the ANR project INSTINCT (N° ANR-22-CE08-0012-01). We acknowledge SOLEIL for provision of synchrotron radiation facilities, and we would like to thank C. Mocuta for assistance in using beamline DIFFABS (proposal 20230095).

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