Design and Fabrication of Highly Elastic and Largely Deformable Nanolaminate Amorphous-Crystalline Metallic Electrodes for Deformable Displays

doi:10.21203/rs.3.rs-4539893/v1

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Design and Fabrication of Highly Elastic and Largely Deformable Nanolaminate Amorphous-Crystalline Metallic Electrodes for Deformable Displays

https://doi.org/10.21203/rs.3.rs-4539893/v1

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The next-generation stretchable displays require a new class of multifunctional electrode materials that simultaneously attain high elasticity, reliable tensile stability, and superior electrical conductivity. Despite their outstanding electrical conductivity, conventional metal thin-films often suffer from limited elasticity, typically less than 1%. Conversely, amorphous alloys offer exceptional elasticity, but their electrical resistivities do not meet industrial requirements. In this work, we propose metallic nanolaminates composed of alternately-stacked Al-based amorphous alloys and nanocrystalline elemental Al layers as a novel kind of electrode material suitable for stretchable displays. These materials are designed to synergistically combine the mechanical properties of amorphous alloys with the electrical properties of crystalline Al, and can be synthesized using currently available mass production fabrication facilities. Through in-situ scanning electron microscopy nano-tension experiments, we demonstrate the nanolaminate samples achieve a unique combination of a high elastic limit approaching 3%, large tensile elongation exceeding 17%, and excellent electrical resistivity lower than 10 µΩ·cm.

Physical sciences/Materials science/Materials for devices/Electronic devices

Physical sciences/Nanoscience and technology/Nanoscale materials/Structural properties

Physical sciences/Materials science/Structural materials/Mechanical properties

Physical sciences/Nanoscience and technology/Nanoscale materials/Electronic properties and materials

stretchable displays

flexible metal electrodes

nanolaminates

metallic glasses

in-situ nano-tension experiments

Prompted by the commercial success of 2D expandable gadgets, such as foldable or rollable displays, recent technological trends in the consumer electronics industry are moving towards the development and commercialization of products operating under much more complex 3D deformations, e.g., wearable, attachable, or multi-directional stretchable screens^1–3. The demand for such devices raises unprecedented challenges in functional materials, requiring them to achieve a new level of multi-functionalities^4–6. For example, while maintaining the traditional functional performances of rigid devices, they must also attain mechanical stability and durability under complex strain conditions. Additionally, compatibility with current facilities and technologies for commercial large-scale production is strongly preferred^3,7.

Currently, two approaches are being pursued to develop highly-stretchable displays. On one side, engineers are exploring the possibilities of building circuits that integrate electrodes, substrates, thin-film transistor (TFT) arrays, and pixels, all utilizing intrinsically deformable materials^8–10. For example, various nanomaterials, such as 1D metallic nanowires (NWs), carbon nanotubes (CNTs), 2D graphenes, MXenes, modified conductive polymers, and liquid metal alloys, are under investigation as potential materials for flexible electrodes^11–15. However, many technical challenges still remain for these materials to be used in the next-generation stretchable displays, i.e., relatively low electrical conductivities compared to traditional metal thin-films, poor mechanical reliability, and complex manufacturing processes. Moreover, synthesis and microfabrication of these materials are not yet compatible with currently available mass production fabrication facilities¹⁶.

An alternative approach utilizes so-called island-bridge architecture, wherein most of the functional components, such as TFTs and pixels, are housed within the rigid islands, and the serpentine-shaped bridges electrically connect them while accommodating the high stretchability of the entire device (see Figure 1a for schematic illustration)^1,17–19. In this architecture, the serpentine bridges consist of dielectric polymer encapsulating parallelly-running conductive metallic electrodes, each called track (see inset of Figure 1a). Here, the geometry of serpentine bridges and the mechanical properties of the composing materials are the key factors that determine the devices' overall stretchability. Particularly, a narrow bridge hosting a high density of tracks is preferred to ensure both high performance and large stretchability^20–22. Figure 1b, c show the distribution of the maximum principal strain calculated using finite element method (FEM) when the device is stretched by 30% along each of the mutually-perpendicular bi-axial directions. Here, 100 pixel per inch (ppi) of the pixel resolution, 100 µm × 100 µm of the island size, and 20 µm of the bridge width are assumed. Notably, stretching the whole device results in uneven strain distribution in the serpentine bridges, with the maximum strain occurring near the inner radius of the arc. For fully reversible deformation-release cycles, it is imperative to keep the strains at each spot below the elastic limits of constituent materials. Particularly, the metallic materials used for conductive electrodes typically have significantly lower elastic limits than the encapsulating polymers. This low elasticity of the metallic electrodes presents a considerable design constraint because they must run through percolating pathways within the bridges, where strains are lower than their elastic limits. For example, Figure 1c compares the regions in the serpentine bridge where the maximum principal strains are below 3.0% (top left) and 1.0% (bottom left) under the 30% of bi-axial device stretching. In other words, these are the zones suitable for metallic electrodes if their elastic limits are either 1.0% or 3.0%. The plot on the right displays the variation of the maximum principal strains calculated along the direction from the inner arc to the outer (see the arrows on the left images in Figure 1c). As is evident, the permissible width for metallic electrodes in the 20-micron-wide serpentine bridge increases by a factor of 8 (from 1.6 to 12.8 μm) as their elastic limit rises from 1.0% to 3.0%. This simple example demonstrates the importance of developing a new electrode material that attains high elasticity without sacrificing electrical properties.

Thin-film metals have been commonly used as electrodes in early-generation deformable electronics due to their high electrical conductivity (>10⁵ S/cm) and manufacturability on flexible substrates. Furthermore, these conventional electrodes offer additional advantages as they can be seamlessly fabricated using existing large-area production facilities built upon conventional lithography and etching processes²³. Nonetheless, they still face significant challenges when aiming to be used in the next-generation stretchable displays. Their elastic strain limits are typically less than 1%^24–26, which serves as a major impediment to developing highly-stretchable electronics.

Amorphous alloys, often called metallic glasses (MGs), typically exhibit extraordinary mechanical properties, such as high strength and large elastic limit^27–29, and have been rising as promising candidates for electrode materials in variable deformable devices^30–32. However, despite the superior mechanical properties, the electrical conductivity of the monolithic MGs is typically an order of magnitude lower than that of the pure metals owing to the presence of additive alloy elements. These are essential to increase the glass-forming ability but contribute to electron scattering. Consequently, MGs fall short of meeting industrial criteria (electrical resistivity lower than 10 μΩ·cm) on their own, limiting their applications in commercial electronic devices.

One effective approach to overcome this trade-off between mechanical and electrical properties is to design nano-scaled composite materials that synergistically combine the large elastic deformability of MGs and the excellent electrical conductivity of elemental metals. The key design challenge here lies in figuring out the spatial configurations of MGs and elemental metals to simultaneously optimize these multiple properties. A large tensile ductility in these materials is additionally desirable to guarantee mechanical reliability during operation. A synthesis route that can utilize the existing large-area production facilities would also offer an extra advantage. Considering all the aforementioned requirements and preferences for electrode materials in stretchable displays, we propose the nanolaminate structure consisting of alternating layers of MGs and crystalline elemental metals^33,34. Here, it is crucial to manage the stacking direction of nanolaminates perpendicular to the directions of mechanical straining and electron transport. Under these conditions, the stronger but less conductive MGs bear a greater load relative to its volumetric fraction³⁵, while the weaker but more conductive crystalline metals provides the fast electron transport channels.

From the perspective of commercial production, it is important to note the limited options for the metallic elements available for conductive electrodes. Although Cu, Ag, and Au are commonly considered potential candidates, Au is typically deemed unsuitable for displays due to its high cost and potential contamination issues during processes. Difficulties in precisely controlling the etching of Ag and Cu would pose technical hurdles for fine patterning. On the other hand, owing to its relatively low cost, high electrical conductivity, and precise etching capabilities, Al is being widely used in the display industry. The low elastic modulus of pure Al offers an additional advantage for its applications under large deformations.

In this paper, we propose the utilization of MG-nanocrystalline nanolaminates as electrode materials that attain both high elastic flexibility and electrical conductivity necessary for next-generation highly-stretchable displays. The MG layers consist of Al-based amorphous alloys (see Table 1 for compositions), which are either purely monolithic or contain tiny nanocrystalline particles. The crystalline layers are made from the elemental Al. The thickness of each layer is maintained between ~30 nm and ~70 nm chosen to optimize both mechanical strength and electrical conductivity simultaneously. The nanolaminate samples comprise a total number of 11 layers, with total thicknesses of around 600 nm. Through the in-situ SEM nano-tension experiments, we demonstrate that these nanolaminate materials can simultaneously achieve elastic limits approaching 3% and electrical resistance lower than 10 μΩ·cm. Additionally, these nanolaminate samples show tensile elongations exceeding 17% or more, indicating excellent mechanical stability under large deformation. The exceptional mechanical behavior of the nanolaminates is further elucidated through molecular dynamics simulations.

Table 1. Composition, geometry, and mechanical properties of the fabricated single-layer amorphous and nanolaminate samples. *These amorphous layers are not fully monolithic but contain tiny nanocrystalline precipitates embedded in the glassy matrix.

Sample name	Structure	Constituent layers	Layer composition	Thickness of individual layers (nm)	Total thickness (nm)	Yield strength (MPa)	Yield strain (%)	Total elongation (%)
Al85S	Single-layer	Amorphous (1 layer)	Al85Y8Ni5Co2	-	599	1490	3.24	3.45
Al85M60	Multi-layer	Amorphous (6 layers)	Al85Y8Ni5Co2	60 ± 2	660	104	2.60	19.86
Al85M60	Multi-layer	Nanocrystalline (5 layers)	Al	60 ± 4	660	104	2.60	19.86
Al90S	Single-layer	Amorphous* (1 layer)	Al90Y10	-	585	1180	3.32	17.72
Al90M67	Multi-layer	Amorphous* (6 layers)	Al90Y10	67 ± 3	589	996	3.11	17.57
Al90M67	Multi-layer	Nanocrystalline (5 layers)	Al	38 ± 8	589	996	3.11	17.57
Al90M50	Multi-layer	Amorphous* (6 layers)	Al90Y10	50 ± 4	517	925	2.65	19.82
Al90M50	Multi-layer	Nanocrystalline (5 layers)	Al	43 ± 6	517	925	2.65	19.82
Al90M32	Multi-layer	Amorphous* (6 layers)	Al90Y10	32 ± 3	467	775	2.38	28.25
Al90M32	Multi-layer	Nanocrystalline (5 layers)	Al	54 ± 8	467	775	2.38	28.25
Al94S	Single-layer	Amorphous* (1 layer)	Al94Y6	-	540	1165	3.12	18.86
Al94M52	Multi-layer	Amorphous* (6 layers)	Al94Y6	52 ± 6	569	837	2.70	25.47
Al94M52	Multi-layer	Nanocrystalline (5 layers)	Al	52 ± 10	569	837	2.70	25.47
NC Al	Single-layer	Nanocrystalline (1 layers)	Al	-	620	291	1.31	26.81

Fabrication and Microstructural Characterization of Amorphous-Crystalline Nanolaminate Thin-films

In this study, we synthesized three distinct types of nanolaminates using three different MG phases: quaternary Al₈₅Y₈Ni₅Co₂, binary Al₉₀Y₁₀, and Al₉₄Y₆. Table 1 provides detailed information on the composition and geometry of nanolaminate samples fabricated in this work. Cross-sectional Transmission Electron Microscopy (TEM) bright field images of the single-layered and nanolaminate samples are given in Figure 2. The quaternary alloy of single-layered Al₈₅Y₈Ni₅Co₂ (Al85S) shows a monolithic microstructure of a fully amorphous phase as in Figure 2a, while the binary alloys of single-layered Al₉₀Y₁₀ (Al90S) and Al₉₄Y₆ (Al94S) in Figure 2b, c contain FCC nanocrystalline phases embedded in the glassy matrix (see Supplementary Figure S1). The diffuse peaks observed in the XRD spectrum of the quaternary alloy (Al₈₅Y₈Ni₅Co₂) presented in Supplementary Figure S2 further support the amorphous nature of this material, while relatively narrow peaks in Al₉₀Y₁₀ and Al₉₄Y₆ alloys suggest the existence of nanocrystalline phases in the binary alloys. In Figure 2a, the amorphous layers of Al₈₅Y₈Ni₅Co₂ in the nanolaminate structure retain fully monolithic amorphousness analogous to the single-layer sample, indicating the lamination process does not alter the microstructure. Similarly, the amorphous layers of Al₉₀Y₁₀and Al₉₄Y₆ alloys in the laminates also contain the nanocrystalline particles embedded within the glassy matrix as they do in the single-layered samples, as shown in Figure 2b, c. Due to the fully amorphous nature, inter-layer interfaces in Al85M60 nanolaminate samples appear smooth. However, those in Al90M67, Al90M50, Al90M32, and Al94M52 samples, all of which contain amorphous layers embedding nanocrystalline precipitates, show relatively large undulations. Figure 2d shows high-resolution transmission electron microscope (HRTEM) images of the amorphous and crystalline layers in nanolaminates together with the Fast Fourier Transform (FFT) patterns in the insets. The size of crystalline particles in Al₉₀Y₁₀and Al₉₄Y₆ alloys is approximately 10 to 20 nm, as in Figure 2d. The nanocrystalline Al layers exhibit a columnar grain structure with the (111) plane oriented perpendicular to the growth direction, and the width of the columnar grains measured as 33 ± 11 nm. Diffuse ring patterns in the insets of the amorphous layers further support the non-crystalline structure, while the sharp spots demonstrate the existence of the crystalline lattice fringes.

Mechanical Characterizations of Nanolaminate Thin-Films

Figures 3 presents results of in-situ nanoscale tensile experiments conducted on the single-layer amorphous and multi-layer nanolaminate samples. In Figure 3a–c, the engineering stress-strain curves for Al₈₅Y₈Ni₅Co₂-, Al₉₀Y₁₀-, and Al₉₄Y₆-based specimens are given. Red curves in each plot represent the stress-strain data of pure nanocrystalline aluminum. Figure 3d–f compares the yield strength and yield strain of each specimen. The specific values of yield strength, yield strain, and total elongation are presented in Table 1. Figure 3g-j shows the representative scanning electron microscope (SEM) images taken during tensile tests for some selected samples, i.e., Al85S, Al85M60, Al90M32, and nanocrystalline Al. Representative in-situ videos of the nano-tension experiments for all samples are available in Supplementary Movies S1-9. As can be seen in Figure 3a, the monolithic Al85S specimens exhibit characteristics of brittle fracture, i.e., linear elastic loading followed by a catastrophic failure. In those samples, no notable evidence of necking is observed, as shown in Figure 3g. Instead, an abrupt shear fracture, facilitated by highly-localized shear deformation, takes place immediately after the elastic regime terminates at 3.2% strain. This brittle-fracture-like catastrophic failure is a typical trait of many metallic glasses³³. In contrast, the pure nanocrystalline aluminum (NC Al) specimen shows a significant ductile elongation up to ~26.8% and necking, as shown in Figure 3a, j. However, the yield strength is considerably lower than that of the monolithic Al85S specimens. Moreover, their yield strain of 1.31%, representing the strain at yield strength and almost equivalent to the elastic limit, is also considerably lower than that of Al85S, which is 3.24%. On the other hand, the nanolaminate sample, Al85M60, has lower yield strength compared to that of Al85S, while total elongation significantly increased from ~3.5% to ~19.9%. Its yield strain also increases by a factor of 2 from 1.31% to 2.60%. Failure mode of the Al85M60 is a ductile fracture, as opposed to the brittle-fracture-like catastrophic failure observed in Al85S specimens. The cup-and-cone type contour appearing in the region near the fracture surface in Figure 3h further demonstrates the ductile nature of these samples. Of particular interest is the SEM image of the Al85M60 specimen in Figure 3h, captured at 13.3% elongation, which reveals the formation of double necking (indicated by red arrows). This observation qualitative implies that this material exhibits excellent resistance against tensile instability.

Unlike the Al85S samples, both Al90S and Al94S single-layer specimens exhibit a considerable amount of plasticity even in the absence of the ductile nanocrystalline Al layers as shown in Figure 3b, c. As is presented in TEM images in Figure 2, these samples are not monolithic metallic glasses but contain crystalline precipitates embedded in the amorphous matrix. The extended plasticity in Al90S and Al94S samples is presumably attributed to the existence of such crystalline phases, which may impede the rapid propagation of shear bands^36–38. Analogously to Al₈₅Y₈Ni₅Co₂-based nanolaminates, the yield strain of samples built with Al₉₀Y₁₀ or Al₉₄Y₆ amorphous layers increases with the volume fraction of nanocrystalline Al layers, approaching or exceeding 3.0%, respectively, as shown in Figure 3e, f. In these samples, the total elongation also scales with the volume fraction of nanocrystalline Al. Especially, Al90M32 exhibits a significantly high total elongation of ~28.2%, even surpassing that of the pure nanocrystalline Al, as shown in Figure 3i. According to the theory for mechanical behavior of composite materials³⁵, strains imposed on each phase in the nanolaminate samples are equal when the tensile loading axis is perpendicular to the stacking direction of layers. Consequently, the strength of the nanolaminates, s_comp, becomes the volume-weighted average of strengths of the individual phases³⁵: s_comp = V_amor s_amor + V_Xtal s_Xtal, where V_amor and V_Xtal are the volume fraction of amorphous and nanocrystalline Al phases, respectively and, s_amor and s_Xtal denote strengths of individual amorphous and nanocrystalline Al phases, respectively. The dashed lines in Figure 3d–f represent the yield strengths of the nanolaminate materials predicted by the conventional theory mentioned above. Noticeably, the actual yield strengths of nanolaminates measured in this study exceed the theoretical estimations. In order to investigate the physical origin of these increased strengths beyond the conventional limits and to elucidate the specific deformation mechanism associated with these unique phenomena, we conducted further analysis using the computational method.

Molecular Dynamics (MD) simulations were performed to investigate the origins of enhanced strengths observed during tensile experiments and the associated plasticity mechanism of nanolaminates composed of Al-based metallic glass and nanocrystalline Al layers. We prepared four different types of specimens, all with the same sample size of 20×20×100 nm³ but different layer configurations: a single-layer columnar nanocrystalline Al, a multilayer nanocrystalline Al, and two multilayer nanolaminate samples with different phases placed on surfaces. The thickness of an individual layer in nanolaminate samples is 4 nm. The first column of Figure 4 illustrates these sample structures. The purpose of the two different stacking sequences in Figure 4c, d is to figure out the effect of the surface layer on the mechanical behavior of nanolaminate samples. In the absence of pre-existing dislocations, the plasticity in the single-layer nanocrystalline Al shown in Figure 4a commences with nucleation of dislocations at grain boundaries (GBs), followed by their gliding to free surfaces. Particularly, locations where GBs intersect free surfaces act as sites for extensive plastic deformation, eventually leading to the development of embryonic cracks (supplementary Movie S10). Ultimately, cracks begin to grow, causing plastic localization and failure due to plastic instability. The addition of Al/Al interfaces, as in Figure 4b, appears to temporarily mitigate the localized plasticity during the initial stage of deformation (compare the atomic shear strain distributions in Figure 4a, b) because inter-layer interfaces block the passage of dislocations. Nonetheless, the plastic instability eventually initiates on the surface and propagates across the layers by forming severe shear bands (supplementary Movie S11). On the contrary, the deformation of the nanolaminate sample shown in Figure 4c exhibits distinct deformation characteristics compared to those observed in the nanocrystalline Al. Although the initial plasticity begins with the dislocation nucleation within the nanocrystalline layer, similarly to the nanocrystalline Al samples in Figure 4a, b, the amorphous layers hinder shear deformation from directly traversing across the adjacent layers. Instead, plastic strains considerably disperse in the amorphous layer (supplementary Movie S12). As a result, each nanocrystalline Al layer in the nanolaminate sample deforms somewhat independently, mitigating the formation of severe shear localization that runs across the entire sample. As the tensile elongation proceeds further, local thinning of the amorphous layers becomes evident in a region close to the highly deformed nanocrystalline Al layers, which eventually evolves into plastic instability and final failure (supplementary Movie S12). As shown in the insets of Figure 4c, this instability, manifested by the embryonic crack nucleation, initially forms in the outmost nanocrystalline layers where GBs meet free surfaces and begin growing with further tensile elongation. The susceptibility of crack initiation at GB sites in the surface nanocrystalline layers suggests that the resistance against failure would improve by physically protecting these vulnerable locations. To clarify the effect of surface condition on plasticity, we changed the layer sequences, positioning the amorphous layers on the outmost surfaces. Simulation results presented in Figure 4d illustrate that the formation of the embryonic cracks is effectively suppressed when GB-surface intersections are eliminated, leading to a substantial alleviation of strain localization. This dispersed plasticity contributes to enhanced tensile ductility, allowing the formation of multiple necking (Figure 4d and supplementary Movie S13). Indeed, the snapshot image in Figure 3h captured during the in-situ tests visually demonstrates the occurrence of the double necking in the actual tensile experiments. In addition, it should be noted that due to substantial differences in spatiotemporal scales, a direct comparison of mechanical properties from the MD simulations with those from the nano-tension experiments is inappropriate. Instead, the results from the MD simulations should be used as a qualitative guideline to understand the differences in deformation mechanisms among different single-layer and nanolaminate samples.

Electrical Characterizations of Nanolaminate Thin-films

Figure 5 shows the electrical resistivities of single-layer amorphous and 11-layered nanolaminate samples as a function of the volume fraction of the amorphous phases. The monolithic metallic glass of Al85Y8Ni5Co2 has a considerably high electrical resistivity of 68.5 μΩ·cm, while those of the other two single-layered amorphous alloys, i.e., Al90Y10 and Al94Y6, are 22.9 and 17.7 μΩ·cm, respectively. Nonetheless, the electrical properties of single-layered amorphous alloys do not meet industrial criteria for electrodes in flexible devices, which typically require electrical resistivity lower than 10 μΩ·cm. On the other hand, the electrical resistivity drastically decreases close to or below 10 μΩ·cm when the nanocrystalline Al layers are added.

In general, thin-film metals exhibit higher resistivities compared to their bulk counterparts³⁹. This increase arises due to electron scattering at surfaces of thin-films or grain boundaries, which is non-negligible when the film thickness or grain size becomes comparable to the electron mean free path (EMFP). According to classical theories by Fuchs and Sondheimer (FS)⁴⁰ and Mayadas and Shatzkes (MS)⁴¹ for electron scattering at external surfaces and grain boundaries, respectively, an extra term proportional to ρ_o𝜆/d, should be added to the resistivity of thin-films, where ρ_o and 𝜆 are the bulk resistivity and mean free path for electron-phonon scattering and d is the length scale representing extrinsic or microstructural dimensions, e.g., film thickness or grain size⁴². The thicknesses of nanocrystalline Al layers in the nanolaminate samples range from ~40 nm (Al90M67) to ~60 nm (Al85M60), sufficiently exceeding the EMFP of bulk elemental Al, estimated at approximately 20 nm⁴². Accordingly, 𝜆/d ratio of nanolaminate samples in this study varies between ~0.33 and ~0.5, suggesting that the thickness effect on the resistivity of nanocrystalline Al layers remains sufficiently small.

The above semi-quantitative estimation of the size effect on electrical properties can be further supported by comparing the experimentally measured resistivities to the theoretical predictions from the parallel-resistors model, in which the resistivity of each layer in the nanolaminates is assumed to remain constant regardless of the layer thickness. In this case, the resistivity of a whole nanolaminate sample becomes:

, where ρ_amor and ρ_Xtal are resistivities of amorphous and nanocrystalline Al layers at room temperature, respectively, and t_amor and t_Xtal are the thicknesses of individual amorphous and nanocrystalline layers, respectively. N_amor and N_Xtal are the total numbers of amorphous and nanocrystalline layers, respectively. The dashed lines in Figure 5 indicate predictions from the parallel-resistors models obtained by substituting ρ_Xtal and ρ_amor, which were measured independently for the thick single-layer nanocrystalline Al and amorphous alloy, respectively. The deviation of the experimental data from this model seems to be minimal. Given that small undulations of the amorphous-nanocrystalline interfaces and presence of grain boundaries in the nanocrystalline Al layers could introduce additional effects, this deviation from the model appears reasonable.

This work presents the development of a novel class of electrode materials that simultaneously attain the excellent mechanical performances of metallic glasses and electrical properties of nanocrystalline elemental metals. The nanolaminate samples in this work achieve a unique combination of a high elastic limit approaching 3%, large tensile elongation exceeding 17%, and excellent electrical resistivity lower than 10 µΩ·cm.

Preparation of the Al-based metallic glass alloy single target

Three types of the Al-based metallic glass alloy single targets with a nominal atomic composition of Al₈₅Y₈Ni₅Co₂, Al₉₀Y₁₀ and Al₉₄Y₆ (at%) were fabricated by powder metallurgy process. Master alloy ingots were prepared through the vacuum plasma melting system (Eloi Materials, Korea) with high-purity elements (over 99.9%) in an argon atmosphere. All ingots were remelted several times to ensure compositional homogeneity. Each powder was fabricated from the master alloy ingots using an argon gas atomizer, which was then sintered by hot isostatic pressure (HIP) at a pressure of 60 kg/mm² and a temperature of 813 K to produce consolidated billet. After cutting and polishing, the fabricated disk- shaped targets had a diameter of 127 mm and a height of 6.35 mm.

Nanolaminate thin-film deposition

The nanolaminate thin-films were prepared with alternating layers of Al-based metallic glass and crystalline pure-Al using DC magnetron sputtering system (Supplementary Figure S4). Si (100) wafers with a diameter of 100 mm were used as substrates after cleaning them in ethanol via ultrasonication before deposition. Base pressure was 2.0 × 10^–5 Torr or lower and a working pressure of 1.0 × 10^–3 Torr was maintained with a pure Ar gas flow rate of 20 sccm. Deposition was carried out with sputtering power of 300 W applied on the targets at room temperature, and the distance between the target and the substrate was maintained at a constant of 120 mm. The cooling channel was inserted at the bottom of the substrate holder and the temperature rise of the substrate was suppressed by circulating coolant at a temperature of 15 ℃ or below. The microstructure of the as-deposited thin-films and thickness of each layer were observed by using transmission electron microscopy (JEM-2100F, JEOL Ltd., Japan). The crystal structures were evaluated using X-ray diffraction (D8-Discover, Brucker, Germany), and the average composition of the thin-films was confirmed by energy dispersive X-ray spectroscopy (X-Max^N, Oxford Instruments, UK).

Measurements of electrical resistivity

Sheet resistance was measured with four-point probe sheet resistance meter (MCP-T610, Mitsubishi Chemical, Japan), and measured sheet resistance values were cross-checked also using a multimeter. The single-layer and 11-layered nanolaminate thin-films were deposited onto SiO₂ coated Si (100) wafer with dimensions of 50 mm × 50 mm. Based on the size of the samples, the measured sheet resistance was corrected with a correction factor of 4.5324, and the electrical resistivity was calculated by multiplying the thickness of the fabricated thin-films.

Nano-tension Experiments

Uniaxial tension tests were conducted on both single-layer and nanolaminate samples. The tensile direction of the dog-bone-shaped tensile specimen is aligned in-plane with the films. The cross-section of the nanopillars has a square shape, with each edge rotated by 45° to the stacking direction, and the length of each side is 320 ± 20 nm. Within the cross-section, there are totally ~ 8 layers included (Supplementary Figure S5). The length of the gauge section is 1500 ± 20 nm, so the aspect ratio of the tensile specimens is 1 to 5. Tensile specimens were fabricated using the focused ion beam (Helios G5, Thermo Fisher Scientific) milled with final ion beam conditions of 30 kV accelerating voltage and 7 pA beam current. In-situ tensile tests were performed using the Picoindenter (PI- 87, Bruker) equipped with custom-made tensile grips in the scanning electron microscope (SU5000, Hitachi). The prescribed strain rate was 1 × 10^− 3 s^− 1 for all samples, and the specimens were pulled until they fractured. To cross-check the accuracy of displacement measurement by the Picoindenter, we visually monitored the actual length change of the gauge sections directly in the movies recorded during the in-situ tensile experiments.

MD simulations

Simulation samples were prepared by assembling separately-created NC Al and MG layers using Matlab code. The NC Al layers were made by the Voronoi tessellation method implemented in Atomsk software⁴³, ensuring that all grains exhibit a < 111 > textured crystal orientation representing the actual columnar structure of the nanocrystalline Al layer (Supplementary Figure S6). The MG layers were also built based on an Al₇₀Ni₂₀Co₁₀ amorphous metallic alloy, with each layer having a thickness of 4 nm. The final samples were then assembled by alternating NC Al and MG layers along the stacking direction with dimensions of 100 × 20 × 20 nm³. The number of atoms in constructed samples varies between 2,424,000 and 2,730,000 atoms based on their layer compositions. MD simulations were then performed using the LAMMPS software⁴⁴ at 300 K with free surface boundary conditions along all directions. Before the mechanical testing, samples were heated to 300 K and equilibrated for 50 ps with the Nosé-Hoover canonical (NVT) ensemble. The uniaxial tensile load was then applied along the axial direction of the specimen by fixing one side of the sample and moving the other in the outward direction with a constant strain rate of 2 × 10⁹ s^− 1 up to 50% total strain. We used the common neighbor analysis (CNA) method^45,46 to visualize the local structure of each atom and atomic strain analysis⁴⁷ to determine atomic shear strain magnitude. The embedded atom method (EAM) potential developed by Pun et al.⁴⁸ was utilized to describe atomic interaction in the Al-Ni-Co system.

FEM simulations

Mechanical simulations of the stretchable pixel array with serpentine bridges were performed using commercial software ABAQUS (Dassult Systemes, France). The FEM models were designed to analyze strain distribution for the in-plane deformation of the multilayers including PI substrate, pixel circuits, Al-metal electrodes, and inter-metal dielectric layers. For computational efficiency, the vertical layout of the FEM model was simplified by replacing Young’s modulus and Poisson’s ratio of all layers with the equivalent properties, except for the electrodes. All layers were assumed to be linear elastic, and 8-node hexahedral brick solid elements were used with the reduced integration method (C3D8R).

COMPETING INTERESTS

The authors declare no competing interests.

AUTHOR CONTRIBUTIONS

W.S. and D.A. contributed equally to this work. D.J. and Y.K. proposed and supervised the project. W.S. and T.H.L. designed the study and developed fabrication methods. D.A. conducted nano-tension experiments. H.G. conducted MD simulations. E.S.P. and K.P. supported experiments. W.S., D.A., H.G., and T.H.L. wrote the manuscript. D.J. revised the manuscript. All authors read and approved the manuscript.

ACKNOWLEDGEMENTS

This work was supported by the National Research Foundation of Korea (NRF) grant funded by the Korean government (MSIT) (NRF- 2020M3D1A1110522, RS-2024-00339298).

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Design and Fabrication of Highly Elastic and Largely Deformable Nanolaminate Amorphous-Crystalline Metallic Electrodes for Deformable Displays

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Abstract

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Introduction

Fabrication and Microstructural Characterization of Amorphous-Crystalline Nanolaminate Thin-films

Mechanical Characterizations of Nanolaminate Thin-Films

Electrical Characterizations of Nanolaminate Thin-films

Conclusion

METHODS

Preparation of the Al-based metallic glass alloy single target

Nanolaminate thin-film deposition

Measurements of electrical resistivity

Nano-tension Experiments

MD simulations

FEM simulations

Declarations

COMPETING INTERESTS

AUTHOR CONTRIBUTIONS

ACKNOWLEDGEMENTS

References

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Supplementary Files

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