# Impact of the brittle layers on the electrical and mechanical behaviours of 18 nm Ag layers in indium tin oxide-silver stackings

Thibault Chommaux, Pierre Godard, Dominique Thiaudière, Pierre-Olivier Renault

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Thibault Chommaux, Pierre Godard, Dominique Thiaudière, Pierre-Olivier Renault. Impact of the brittle layers on the electrical and mechanical behaviours of 18 nm Ag layers in indium tin oxide-silver stackings. Thin Solid Films, 2025, 826, pp.140765. 10.1016/j.tsf.2025.140765. hal-05387992

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# Impact of the brittle layers on the electrical and mechanical behaviours of 18 nm Ag layers in indium tin oxide-silver stackings

T. Chommaux<sup>a,b</sup>, P. Godard<sup>a,*</sup>, D. Thiaudière<sup>b</sup>, P.O. Renault<sup>a</sup>

## Abstract

The electrical and mechanical properties of indium tin oxide (ITO) and silver (Ag) multilayers deposited on polyimides (PI) are studied. From the electrical point of view, the multilayer performs initially like the Ag layer, as seen in the *R*-sheet and in the gauge factor, and behaves during biaxial tensile tests according to the mechanical properties of this 18 nm Ag layer. The mechanical properties strongly change for different multilayer architectures. For example, the ITO/Ag/ITO/PI sample shows an elastic regime, whereas plasticity occurs immediately in Ag/ITO/PI, ITO/Ag/PI and Ag/PI samples. Moreover, the Ag layer sustains a stress of almost 500 MPa before cracking in the first case, but no more than 200 MPa in the second. This shows how the mechanical behaviour of ultra-thin silver layers is dictated by the interfaces, that may be soft (as with air or polyimide) or hard (amorphous or crystalline).

**Keywords:** Thin films; Multilayers; Fracture; X-ray diffraction; Electrical resistivity

## 1. Introduction

Indium Tin Oxide, or ITO, is the most frequently used transparent conducting oxide. The crystalline phase shows improved optical and electrical properties as compared to the amorphous phase [1]. Yet, sputtered ITO films deposited on compliant substrates are mainly amorphous. Several strategies have been adopted to increase the crystallinity of the films, like surface treatments of the substrate [2], annealings [3, 4] or the deposition of a crystalline interface between the ITO layer and the substrate. The latter may have other benefits, e.g. if it is a ductile metal. To avoid optical reflection, this metal layer is often encapsulated between the ITO layers [5, 6, 3, 7, 8].

During manufacture or operation, multilayer films on flexible substrates undergo complex stress distributions that can result in plastic deformation or brittle fracture. In a bi-or tri-layer architecture, the combination of a ductile material with a brittle material can improve the mechanical properties. The mechanical behaviour of bi- and tri-layers of various metals have been successfully studied under quasi-equi-biaxial mechanical loading with X-ray diffraction (XRD) to provide a new understanding of the influence of layer position [9]. If the microstructure of a ductile metallic material reaches nanometre dimensions, a ductile-brittle transition can take place because of the high interface density or the reduced thickness and grain size [10, 11, 12].

This study focuses on the electro-mechanical properties of ITO/Ag/ITO multilayers deposited on a polymeric substrate during tensile tests. Not only is the resistance measured, but also the stress in each layer thanks to XRD. To have sufficient signal-over-noise ratio in these in-situ measurements, we had recourse to a synchrotron radiation facility. The use of XRD provides the stress evolution in each layer simultaneously and allows understanding the damage mechanisms occurring in the different crystalline stackings. To this end, five deposits are compared: noting the layers from the top to the bottom, we considered the trilayer ITO/Ag/ITO/PI, two bilayers Ag/ITO/PI and ITO/Ag/PI, and two monolayers ITO/PI and Ag/PI, where PI stands for the polyimide substrate (125 µm thick Kapton®, not counted as a layer). The Ag layer thickness was 18 nm; though most authors find that ITO/metal/ITO samples perform better when the metallic layer is slightly thinner [13, 14, 8, 15], some authors showed that 18 nm corresponds to stacking with a very high figure of merit [16]. Moreover, the tensile tests are equibiaxial. This is generally closer to the service conditions than uniaxial tests, and avoids buckling [17]. Furthermore, in our case, it allows the sample to remain perfectly flat during the mechanical test. Comparisons between uniaxial and equibiaxial tests are given in [18, 19]. The 5 different stackings allow for the study of the mechanical behaviour of an ultrathin metallic Ag film in different configurations: confined between two layers of a brittle material (ITO), confined between a brittle material (ITO) and soft material (a polymer or the air), and finally both of its faces exposed to a polymer or the air.

<footnote><sup>a</sup>Université de Poitiers - CNRS - ENSMA, Futuroscope Chasseneuil, France</footnote>

<footnote><sup>b</sup>Synchrotron SOLEIL, L'Orme des Merisiers, Saint-Aubin, Gif-sur-Yvette Cedex, France</footnote>

<footnote>*Corresponding author; address: TSA 41123, 86073 Poitiers Cedex 9, France; email address: pierre.godard@univ-poitiers.fr</footnote>
## 2. Methods

The layers were deposited by argon beam sputtering at ambient temperature in a Nordiko 3000 chamber. A 5 sccm O₂ flux was added during the ITO deposition (In₂O₃/SnO₂ and Ag targets), without acceleration voltage. As shown in [20], this O₂ flux improves the mechanical properties of ITO films, and in particular the crack onset strain. The resulting working pressure was 10⁻⁴ Pa. To avoid the contamination of O₂ inside the Ag layers, a delay of 500 s was imposed between the ITO and the Ag depositions. The ITO and Ag growth rates were 0.13 and 0.28 nm/s, respectively, as determined by reflectometry. This resulted in thicknesses of 70 and 18 nm, necessary to have Ag continuity and a good signal in in-situ X-ray diffraction. The samples' shape are squares of 12 mm side, the width of the cruciform substrate arm is 20 mm.

The experiments were performed at the DiffAbs beamline of the SOLEIL synchrotron radiation facility. For XRD, a beam with an energy of 18.00 keV and a size of ≈ 300 × 250 µm² (horizontal times vertical) was used. The diffraction patterns were collected on an XPAD-S140 two dimensional detector with an exposure time of 4 s. Positioned at 621 mm from the sample, it covers a 2θ range of ≈ 7°. The incident angle of the X-ray beam was fixed to 7° and the planes parallel to the film surface were probed. This low incident angle insures a beam footprint over 2 mm in one direction, and so a high representativity of the XRD signal. The sample resistance, or more properly its electrical R-sheet, was measured with a Keithley 2701 multimeter in a four-point configuration (8 measurements to eliminate the contributions from the wires and possible electrometers offsets). A current of 1 mA was imposed, and the voltage was read every 1.5 s. For single-phase film, the R-sheet can be converted to the resistivity with ρ = Rd, where d is the film thickness.

The tensile machine, locally developed, was presented in [21, 22]. The strain-rate was 3.6 × 10⁻⁶ s⁻¹. A pre-loading of 20 N insures the sample stability and flatness, essential for XRD; this corresponds to an elastic strain of ≈ 0.1 %. Below the machine lie a teleobjective lens and an optical camera (Pixelfly from PCO AG, 1392 × 1024 pixels with a dynamic range of 12 bits and a pitch of 6.45 µm) that allow the measurement of the in-plane strain components εₓₓ and εᵧᵧ imposed to the substrate with Digital Image Correlation, or DIC (Corelli software, [23]). The speckle pattern consists of paint sprayed onto the substrate surface. As assessed in [22], the biaxial tester induces homogeneous strain at a sub-millimetric scale and a displacement uncertainty of about 10⁻² pixels. The samples were stretched up to ≈ 1.2%, i.e. beyond the ITO crack onset strain. The presented strain values are averages over a ≈ 6 × 9 mm² area, where it was homogeneous. Furthermore, the paint contains TiO₂ particles that induce continuous Debye-Scherrer rings on the XRD patterns; this was used as a reference powder.

<table><caption>Table 1: Electrical resistance R-sheet, resistivity ρ (for monolayer samples) and gauge factor G as measured with the Van der Pauw method.</caption><thead><tr><th>sample</th><th>R (Ω/□)</th><th>ρ (Ω·cm)</th><th>G (1)</th></tr></thead><tbody><tr><td>ITO/Ag/ITO/PI</td><td>3.6</td><td>×</td><td>2.5</td></tr><tr><td>Ag/ITO/PI</td><td>5.3</td><td>×</td><td>2.6</td></tr><tr><td>ITO/Ag/PI</td><td>4.4</td><td>×</td><td>2.8</td></tr><tr><td>Ag/PI</td><td>5.2</td><td>9.3 10<sup>-6</sup></td><td>2.8</td></tr><tr><td>ITO/PI</td><td>210</td><td>1.4 10<sup>-3</sup></td><td>-2.7</td></tr></tbody></table>

## 3. Results and discussion

### 3.1. Initial Characterization

Diffactograms at the beginning of the tests are shown in figure 1; one observes the (222) ITO and the (111) Ag Bragg peaks. The absence of {400} or {200} planes indicates strong {111} textures for both layers. The Scherrer formula, with an instrumental contribution of 0.03° to the peaks' FWHM, allows an estimation of the grain size around 13–14 nm in ITO, 16–19 nm in Ag. The ITO/PI film is not shown, because the only observed Bragg peaks correspond to the TiO₂ nanoparticles. Hence, the ITO layer in this film is amorphous; this is confirmed with the absence of ITO peaks in Ag/ITO/PI. From this we deduce that the ITO peak observed in the ITO/Ag/ITO/PI film pertains to the top layer (grown on crystalline silver and not on Kapton). That the top ITO layer shows Bragg peaks whereas the bottom one is amorphous has already been encountered, e.g. in [6]. We note that in reference [20] the ITO was at least partly crystalline, but its thickness was much larger (600 nm) and the sputtering gas was Xe. Moreover, it was shown in this article that the introduction of O₂ at 5 sccm greatly decreases the degree of crystallinity of the ITO films.

The reference positions of stress free ITO and Ag phases are also shown in figure 1 (ICDD® reference). It is obvious that the planes parallel to the surface are largely dilated for the ITO layer, slightly for the Ag one. To be more precise, the sin²ψ method was applied before the mechanical test to determine the residual stress of the layers. The crystalline ITO layers have an in-plane residual strain of −1.90 % measured on the (222) Bragg peak. This corresponds to a residual stress of −2.2 GPa (with E = 116 GPa and ν = 0.35 [24]). We found εₓₓ⁽²²²⁾ = εᵧᵧ⁽²²²⁾ = −0.05 % for the (222) peak of the Ag layer, corresponding to a residual stress of −0.07 GPa (with ν = 0.37 and E = 85 GPa [25]).

The as-deposited R-sheet of the different samples is given in table 1. The comparison between the ITO/PI film and the other ones shows that 18 nm of Ag greatly enhances the electrical conduction in the multilayer films. For single-phase films, the resistivity ρ is also given. The resistivity of bulk Ag is 1.59 × 10⁻⁶ Ω·cm at 300 K [26]. The value here is almost six times larger, due to the grain size, and possibly to other crystalline defects. For ITO, the value is close to the one reported in [20] for a 600 nm
<figure><img src="image_2.png" /><figcaption>Figure 1: Diffractograms of the samples under study with zooms on ITO (222) planes, superimposed in the insert. The dashed lines correspond to stress free crystals. The ITO/PI system is not shown since ITO layers deposited on Kapton are amorphous. The schematics on the right give the systems' architecture.</figcaption></figure>

thick film: $1.5 \times 10^{-3} \Omega \cdot cm$. This is comparable to resistivities found for sputtered films, e.g. as $4.0 \times 10^{-3} \Omega \cdot cm$, (for 100 nm film) [6], $0.6 - 0.7 \times 10^{-3} \Omega \cdot cm$ (100-250 nm) [4], $1.6 \times 10^{-3} \Omega \cdot cm$ (250 nm) [27], $0.9 \times 10^{-3} \Omega \cdot cm$ (200 nm) [2].

## 3.2. In-situ tensile test: electrical behaviour

Mechanical tests were applied to all samples; the results are given as functions of the applied strain $\epsilon_m = (\epsilon_{xx}^{DIC} + \epsilon_{yy}^{DIC})/2$ as measured with DIC. An example of $\epsilon_m$ is given in figure 2. The mechanical anisotropy of the substrate is observed.

The sheet resistance variations at the beginning of the tests are shown in figure 3. It is observed that they first follow a linear evolution, with slopes between 2.5 and 2.8 for the film with an Ag layer, -2.7 for the ITO/PI film. Hence, the Ag layer determines not only the initial R-sheet values in multilayers (as shown in table 1), but also its initial derivative.

<figure><img src="image_3.png" /><figcaption>Figure 2: Macroscopic deformation at the sample centre as obtained from digital image correlation. Nearly equibiaxial tests are applied, and the results are given as functions of the mean values $\epsilon_m$.</figcaption></figure>

<figure><img src="image_4.png" /><figcaption>Figure 3: Zooms on the R-sheet evolution at the beginning of the tests. The initial slopes give the gauge factors.</figcaption></figure>

These slopes at the beginning of the tensile test correspond to the gauge factor $G = (\Delta R/R)/\epsilon_m$ and are also reported in table 1. Repeatability tests gave a dispersion of 0.1. Note that these values are not corrected by the geometric factor (the thickness of the film decreases while it is elongated), but this accounts for less than 10 % of the given values.

Figure 4 shows complete views of the resistance variations during the equibiaxial tests. Large hysteresis loops are observed for samples that contain at least one ITO layer, but only a low (< 5%) and quickly reversible increase in the Ag/PI sample. Upon unloading, the ITO/PI film shows a constant resistance from $\epsilon_m \approx 1.0\%$ to $\approx 0.5\%$. The resistance of the other samples decreases as soon as the straining direction is reversed.
<figure><img src="image_5.png" /><figcaption>Figure 4: R-sheet relative variations (red), elastic strains $\epsilon_{zz}^{(222)}$ of ITO (grey) and $\epsilon_{zz}^{(111)}$ of Ag (green) as functions of the mean applied strain $\epsilon_m$ for the five samples. (a) ITO/Ag/ITO/PI, (b) Ag/ITO/PI, (c) ITO/Ag/PI, (d) Ag/PI and (e) ITO/PI. The loading (resp. unloading) is represented in full (resp. open) symbols. The dashed lines correspond to a fully elastic strain transmission from the substrate to the film; they are different for the two layers because of the elastic anisotropy of silver.</figcaption></figure>

### 3.3. In-situ tensile test: mechanical behaviour

Together with the R-sheet evolution, figure 4 shows the ITO and Ag elastic strains $\epsilon^{(hkl)}$ in the crystalline layers. The dashed lines correspond to the slopes $s_{ITO}$ and $s_{Ag}$ that would be obtained if the macroscopic strain imposed to the substrate was fully transmitted as an elastic strain in the layers (gray for ITO, green for Ag); they are materials dependent because of the elastic anisotropy (Ag has a large Zener anisotropy ratio of approximately 3, whereas we showed in [20] that ITO has a very low elastic anisotropy, confirming results from [28]). The ITO layers follow first a linear regime with a slope equal to $s_{ITO}$, evidencing an elastic regime and a full adhesion to the substrate. This is seen in the two samples with a crystalline ITO layers, that is ITO/Ag/ITO/PI and ITO/Ag/PI. Depending on the stacking, the Ag layer has, for the ITO/Ag/ITO/PI sample only, an $\epsilon^{(111)}$ vs $\epsilon_m$ slope equal to $s_{Ag}$ at the beginning of the test; for the other samples, the elastic strain curve is much lower (in modulus) than the $s_{Ag}$ curve. However, whatever the case, there is a full strain transmission from the substrate to the top layer of the thin film. Indeed, the full strain transmission is evidenced by the ITO layers above Ag, that can have an elastic strain following $s_{ITO}$ in the linear regime (figure 4a and c) only if the layers below fully transfer the strain imposed by the substrate - as an aside, this shows that the strain transfer is also complete through the amorphous ITO layer in the ITO/Ag/ITO/PI sample. The full strain transmission from the substrate to the film and between different layers of a film was also evidenced in [22, 29]. We note also that the Ag layers are in a compressive state after the unloading. This is attributed to a plastic elongation during the load, and the films go to a negative stress during the unload because they are carried away by the substrate that has a larger elastic limit.

This strain transmission is attributed to the sample elaboration method. Indeed, ion beam sputtering is a high energetic process, leading to adherent and dense films, but containing a high density of crystalline defects. On the one hand, these defects are the source of residual stresses that facilitates the adhesion of the films on the substrates, but on the other hand, these crystalline defects are loci where plastic events appear easily, hence the small yield stress observed in the Ag/PI sample.

With Hooke's law, the elastic strains are converted to in-plane stress in figure 5. Several transitions are observed in the mechanical behaviours of these samples. The crystalline ITO layers are clearly brittle. After an elastic regime (domain E in figure 5a), a second slope is observed during crack nucleation and multiplication (domain C). Indeed, in domain C, the stress increases at a slow rate because the stress relaxation at the vicinity of the cracks partly counterbalances the increase in the areas far enough from any cracks [30]. There are also two different slopes during the unload (see below for the interpretation).

The mechanical properties of the Ag layer depend on its environment, figure 5b and c. On the one hand, when
<figure><img src="image_6.png" /><figcaption>Figure 5: Applied stress $\sigma$ as a function of the macroscopic strain $\varepsilon_m$, as measured from (a) (222) planes in ITO, (b,c) and (111) planes in Ag. Subfigure (b) shows that the mechanical domains in the Ag layer are completely different when its interfaces are soft (air and polyimide) or strong (amorphous and crystalline ITO). The loading (resp. unloading) is represented in full (resp. open) symbols.</figcaption></figure>

isolated (Ag/PI sample), it shows a ductile behaviour, with an elasto-plastic regime EP starting from the beginning of the test, followed by a fully plastic regime P at $\varepsilon_m \approx 0.35 \%$. On the other hand, when sandwiched between ITO layers, it shows three distinct domains during the loading, corresponding to an elastic regime, an elasto-plastic one and a cracking one (domains E, EP and C, respectively, in figure 5b). This will be further discussed in subsection 3.5.

### 3.4. Cracking phenomena

In supported thin films, a crack may not immediately propagates through the sample thickness and/or width. In other words, the sample may simultaneously be locally in regime C and locally in regime EP. X-ray diffraction, measuring the stress averaged over the illuminated area, gives only the sum of two contributions: the stress decrease at the vicinity of the cracks, and the increase in the areas far enough from any cracks [30, 31, 32, 33]. Electrical measurements, on the other hand, give the conduction channels, here between four points situated at the vertices of the squared samples. The channels are slightly disturbed by elasto-plastic phenomena (as seen in the Ag/PI sample), but broken as soon as a crack propagates through the whole film. These two complementary methods are used in the sequel to determine the Crack Onset Strain (COS) $\varepsilon_c$ appearing during the loading, and the possible closure of the cracks during the unloading. In addition, information about the transmission through the different layers of the composites are obtained.

Due to its brittleness, crystalline ITO presents a clear domain E to domain C transition in $\varepsilon^{(222)}$ that allows the determination of the COS $\varepsilon_c$ with X-ray diffraction, table 2. We obtain $\varepsilon_c = 1.00 \pm 0.05 \%$ for both samples ITO/Ag/ITO/PI and ITO/Ag/PI. For the Ag/ITO/sample, there is no ITO Bragg peaks, and the reported value is associated to the transition from the domain EP to the domain C in Ag; the COS is also close to 1.00 %. Of course, $\varepsilon_c$ cannot be determined in this way in the amorphous ITO/PI sample. Finally, there is no domain C in the Ag/PI film.

It is interesting to observe that in samples having two crystalline layers the XRD gives the same COS in both layers; this is well seen in figures 4a and c for the ITO/Ag/ITO/PI and ITO/Ag/PI samples, where Ag shows a clear EP to C transition. This suggests that the cracks propagate quickly through the whole composite thickness. These results show the benefit of using XRD as a complementary tool to more direct methods such as electrical measurements or optical microscopy.

A 10 % threshold in the R-sheet relative variation is chosen for the COS determination, as in [1]. The values are reported in table 2; we specify that the piezoresistivity of the samples has been taken into account, that is $\varepsilon_c$ is the mean strain for which $\Delta R/R$ deviates from $G\varepsilon_m$ by the given threshold. The measured values are within $1.0 \pm 0.1 \%$ for the Ag/ITO/PI, ITO/Ag/PI and
<table><caption>Table 2: Macroscopic strain where crack nucleation is detected with XRD and with resistivity measurements, the second is given for a 10 % threshold. The uncertainty is evaluated to ±0.02% in each case. NM = impossible to measure. NC = no evidence of cracks.</caption><thead><tr><th>sample</th><th>ε<sub>c</sub> (%) from ε<sup>(hkl)</sup></th><th>ε<sub>c</sub> (%) from R</th></tr></thead><tbody><tr><td>ITO/Ag/ITO/PI</td><td>0.96</td><td>0.70</td></tr><tr><td>Ag/ITO/PI</td><td>1.02</td><td>0.91</td></tr><tr><td>ITO/Ag/PI</td><td>1.02</td><td>1.06</td></tr><tr><td>Ag/PI</td><td>NC</td><td>NC</td></tr><tr><td>ITO/PI</td><td>NM</td><td>0.91</td></tr></tbody></table>

ITO/PI samples; the ITO/Ag/ITO/PI sample shows a smaller value of 0.70 %. Hence, except for the most complex architecture (the ITO/Ag/ITO/PI sample), the COS are very similar, whether determined by XRD or by resistivity. This justifies the 10 % criterion: it does not evidence the very first crack nucleation, but rather gives the regime of crack multiplication and extension. Moreover, in our case, it mitigates the fact that the probes are not localized at the same position: the X-ray beam impinges in the centre of the film while the electric contacts are on the exterior of the film, where the applied strain might be slightly larger.

During the unload, the ITO first closes some cracks in domain *F*, before recovering an elastic regime at around ε<sub>m</sub> ≈ 0.5 %, figure 5a. This is also seen in the *R*-sheet measurement that the resistance goes back to the initial value, figure 4. This crack closure prevents any post mortem characterization of the crack pattern.

### 3.5. The behaviour of the 18 nm Ag layers

The simplest sample containing an Ag layer is Ag/PI. In this sample, no strong increase of ρ is observed, nor any well-defined transition in the stress-strain curve; both remarks indicate that no cracks appear in the film. This may come as a surprise, because ultra-thin films tend to become brittle as the thickness decreases.

The stress-strain curves (figure 5c) show that the Ag layer in the Ag/ITO/PI sample deforms plastically even at small strain. These two samples, Ag/PI and Ag/ITO/PI, are the ones for which the silver layer has no interface with crystalline media. Moreover, we see that the stress σ reaches a higher level in Ag/ITO/PI than in Ag/PI (170 vs 120 MPa), that is when silver is deposited on ITO rather than on Kapton. This suggests that the Ag-ITO interface is a harder interface for dislocations than Ag-PI. This is confirmed when the Ag layer is sandwiched between ITO and PI (σ up to 210 MPa), and finally between two ITO layers, where the stress reaches almost 500 MPa. In fact, it is only in the latter case that the Ag layer shows an elastic regime, up to an applied strain of ε<sub>m</sub> = 0.35 %. It is also the sample with the sharpest EP to C transition.

Here we clearly show that an ultrathin silver layer simply deposited on the polymeric substrate keeps its ductility. Moreover, a plasticity deformation is observed in every

Ag layers. Given the nanometric grain size, grain boundary sliding can be an alternative to dislocation activity. If we had only dislocation-induced plasticity, the stress at the end of the test (after unloading) should be higher than what we observe in figure 5c. Additional experiments would be necessary to determine the contribution of each activated mechanism.

Finally, the current experiment clearly illustrates a constraint effect which refers to the restriction of plastic deformation (whatever its mechanism) in a material due to the presence of adjacent layers with different mechanical properties. It is well-known that the interface between different materials or layers can act as a barrier to dislocation motion [34]; it is extended in this article to the other plasticity mechanisms. In a multilayered composite system, the layers can constrain the deformation of each other and the type of interface plays a crucial role in determining the extent of the constraint effect. Here the ITO layers constrain the deformation of the Ag layer, leading to a delay or an attenuation of the dislocation activity and of the grain boundary sliding, and so to a higher stress state within the Ag layer. This reduction of the plasticity activity promotes a brittle behavior even if in the present experiment the brittle behavior is probably induced by the ITO failure.

## 4. Conclusion

This article clearly shows that the electrical behaviour of ITO and Ag stackings is governed by the silver layer; it does not diverge when Ag is surrounded by soft interfaces (air and polymeric substrate), but shows large hysteresis when Ag has hard interfaces (amorphous or crystalline ITO). In other words, the electrical behaviour of the composite depends on the mechanical properties of the metallic layer, that itself depend on the dielectric-metal interfaces. Given that the films were deposited at room temperature on polymeric substrates, this may have important consequences for the design of flexible optoelectronics systems.

We also showed that there is a full strain transmission from the substrate to the top layer in an ITO/Ag/ITO stacking. The crack onset strain is around 1 % of equibaxial applied strain, as checked with X-ray diffraction and electrical measurements. Moreover, thanks to its phase selectivity and its in-situ capacity, X-ray diffraction evidences that the cracks propagate very rapidly across the crystalline layers.

Finally, we exemplified major differences in the behaviour of 18 nm Ag layers, from a brittle behaviour that supports 500 MPa when sandwiched between ITO layers, to a ductile behaviour with a flow stress of 120 MPa when it is simply deposited on a soft substrate. This illustrates the critical role of the interfaces in multilayer systems. That the metallic layer may be ductile or brittle depending on its environment extends the results of [19], where we showed with XRD and post-mortem microscopy that
600 nm Ni films were mainly brittle when stretched equibiaxially, but ductile when stretched uniaxially.

## Acknowledgments

The authors gratefully acknowledge SOLEIL for beam time allocation and Cristian Mocuta for help during the experiments. This work was partially funded by the French Government programs "Investissement d'Avenir" (LABEX INTERACTIFS, reference ANR-11-LABX-0017-01, EUR INTREE, reference ANR-18-EURE-0010). The région Nouvelle-Aquitaine and the French synchrotron SOLEIL supported this work via the NACRES project.

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