Introduction

High integration is essential for a wide array of cutting-edge flexible electronic devices like wearable/implantable bioelectronics1,2, soft displays/touch screens3,4, and stretchable photovoltaics5. This trend positions two-dimensional (2D) or ultrathin nanomaterials as favorable candidates for multifunctionality and optimized performances6,7,8,9. However, despite their promising potential, the high surface-to-volume ratio and chemical reactivity of most nanomaterials render them vulnerable to environmental factors like temperature variations, humidity, stress, absorption, and radiation10. Therefore, pioneering nanomaterial-based devices are susceptible to ambient degradation and damage, such as organic light-emitting diodes (OLED)11, 2D perovskite solar cells12, and layered GaSe semiconductors13. These challenges cause the increased costs in preparation and their limited reliability in real-world applications14. To tackle this issue, the prevailing strategy is to develop advanced encapsulation to meet fundamental requirements including but not limited to conformal contact for seamless adhesion, low water vapor transmission rate (WVTR), mechanical flexibility and durability, as well as optical properties compatible with functional components13,15,16,17. Nevertheless, existing encapsulation techniques struggle to provide these attributes simultaneously15, and the design of multifunctional encapsulation nanomaterials is still immature18.

Metal oxide nanomaterials, renowned for their excellent tolerance to environmental degradation, are widely used in safeguarding, supporting, and packaging applications19,20,21. High-entropy-oxides (HEOs) further exploit the thermodynamic stability empowered by the high-entropy effect, a concept extended from high-entropy alloys (HEAs) to various high-entropy materials22,23,24,25. Consequently, HEO nanomaterials show exceptional resistance to oxidation and corrosion, making them highly valuable in long-term services26,27,28,29. However, a major challenge for HEOs lies in their inherent brittleness, akin to traditional metal oxides, with a fracture strain typically capped at approximately 1%30,31. This limits their application in scenarios requiring deformability and fracture resistance, such as in stretchable sensors32, flexible batteries33, and force-sensitive piezoelectric devices34. Furthermore, mechanical loadings that might be inconsequential for larger-scale materials can lead to nanomaterial rupture in many practical operations35,36. Enhancing the fracture resistance of HEO nanomaterials is thus a crucial step towards fully realizing their multifunctional benefits.

On the other hand, nature’s ever-evolving processes have long served as a wellspring of inspiration for material design, offering a plethora of biological materials renowned for their exceptional fracture resistance and tensile strength, such as articular cartilage, bone, teeth, nacre, and abalone shell37,38,39. A fundamental toughening principle observed in these materials is the strategic interspersion or interpenetration of two or more nano-phases with markedly different mechanical properties40,41. This precise organization of nanostructures gives rise to hierarchical toughening mechanisms, encompassing phenomena like crack bridging, deflection, and energy dissipation42,43, ultimately bestowing these biological materials with remarkable mechanical properties. Drawing from this natural blueprint, similar strategies have found successful application in the development of hydrogels37, graphene oxides44, and bulk metals or alloys45.

In this work, we present an approach involving the in-situ formation of a 2D HEO scaffold interwoven with nano-sized decomposed polymer chains. The resulting freestanding nanomembranes exhibit exceptional fracture resistance, boasting a ductility nearing 90% and a toughness exceeding 300 MJ/m3, attributed to the dual-phase nanostructure. In addition, the complex chemical compositions of the HEO nanomembrane confer it with a blend of advantageous properties, including strong adhesion to diverse substrates, exceptional optical transparency, and inherent oxidation resistance. This combination of toughness and multifunctionality positions the HEO nanomembrane as a versatile candidate for a myriad of applications in challenging environments. For instance, it can be patterned into ultra-thin and transparent thin-film encapsulation (TFE) for flexible devices, showcasing its potential across a broad spectrum of uses.

Results

Preparation and characterization of HEO scaffold nanomembrane

In this study, freestanding nanomembranes were meticulously fabricated using a technique known as polymer surface buckling enabled exfoliation46 (Fig. 1a). Following the well-established method47, our process commenced with spin-coating a thin layer of polyvinyl alcohol (PVA) hydrogel onto a glass substrate, which was subsequently dehydrated by drying in an oven. A follow-up step involved the deposition using a HEA target with a nominal composition of (CoNi)50(TiZrHf)50 (at. %) onto the dehydrated PVA, leading to the exfoliation of nanomembranes through interfacial fracture induced by the surface buckling of the hydrogel when immersed in water (“Methods” and Supplementary Movie 1). It is worth noting that these nanomembranes possessed a thickness of 30 nm, with in-plane dimensions extending to tens of millimeters, while maintaining an average surface roughness as low as approximately 0.65 nm (Supplementary Fig. 1). The bright-field scanning transmission electron microscopy (BF-STEM) image depicted in Fig. 1b reveals a distinctive dual-phase nanostructure, wherein one phase (the dark phase) act as a 2D scaffold, while the other serves as nano-sized inclusions. Figure 1c showcases the high-resolution TEM (HRTEM) image of the two nano-phases alongside their fast Fourier transformation (FFT) patterns, both displaying an amorphous halo ring. This observation aligns with the selected area diffraction pattern (SADP) of the entire dual-phase nanostructure (the inset in Fig. 1b), which indicates an overall amorphous structure of the nanomembrane. We can manipulate the relative proportion of the two nano-phases by adjusting the spin-coating speed of the hydrogel (Supplementary Fig. 2–4). Furthermore, a comparative nanomembrane of a similar thickness was prepared by sputtering the same HEA target onto a NaCl substrate, which subsequently peeled off by dissolving the NaCl substrate in water, resulting in a freestanding HEA nanomembrane characterized by a uniform single phase (Supplementary Fig. 5). We also synthesized a HEO nanomembrane on the NaCl substrate by preparing a HEO target with the chemical composition of the scaffold determined through subsequent characterization. The HEO nanomembrane exfoliated from the NaCl substrate exhibits a compositionally matched amorphous microstructure (Supplementary Fig. 6), enabling a rigorous comparison with our HEO scaffold nanomembrane. Subsequently, elemental mapping was conducted using energy-dispersive X-ray spectroscopy (EDX) on images obtained through the high-angle annular dark-field STEM (HAADF-STEM) technique, as illustrated in Fig. 1d. This mapping identifies the presence of metallic elements (including Co, Ni, Ti, Zr, and Hf) and non-metallic elements (C and O). Interestingly, our observation revealed that the metallic elements were primarily concentrated within the scaffold phase, while C and O were almost uniformly distributed throughout the entire nanomembrane. These findings are consistent with the Z-contrast observed in the HAADF-STEM image, where the bright contrast of the scaffold phase indicates a predominant composition of heavy elements (e.g., metallic elements), whereas the dark contrast of the inclusion phase suggests a prevalence of lighter elements (e.g., C and O)48.

Fig. 1: Processing and microstructural characterization of the HEO scaffold nanomembrane.
figure 1

a Schematic illustration of the synthesis process and potential application as TFE. b BF-STEM image with the inset SADP. c HRTEM image and the corresponding FFT patterns of the two nano-phases. d HAADF-STEM image and EDX elemental mapping.

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To determine the specific compositions of the dual phases, we conducted X-ray photoelectron spectroscopy (XPS) analyses and atomic force microscopy (AFM) characterizations. Figure 2a–g displays the narrow-scan XPS spectra at different etching durations, complementing the EDX results. The peaks of C 1 s indicative of metallic carbides, C = C, C–C, C–O, and C = O bonds were carefully characterized49 (Fig. 2a). Similarly, the O 1 s spectra unveil the presence of metallic oxides, C = O, and C–O (Fig. 2b). Distinct peaks associated with oxides and pure metals were observed, with negligible carbides being detected (Fig. 2c–g). The presence of adventitious carbon and oxides on the surface in the absence of etching can be attributed to ambient contamination and oxidation50. As the etching duration increases, the prevalent chemical states of carbon are predominantly manifested as C = C, C–C, and C–O bonds, while oxygen primarily exists as metal oxides. The metallic elements within the interior of the nanomembrane can be categorized into two distinct groups. Specifically, Ti, Zr, and Hf predominantly persist in their oxide forms, whereas a portion of Co and Ni exist as pure metals. Through peak fitting of the XPS spectra (Supplementary Figs. 7 and 8), we obtained the atomic ratios of each chemical species in metallic elements, as illustrated in Fig. 2h and Supplementary Fig. 9. Only a minor fraction of pure metals (<20%) and metal carbides (<5%) were detected, with the majority of metallic elements forming their respective oxides. According to the well-established concept22,25, the determined chemical composition of the oxides (Co1.6Ni0.7Ti2.1Zr2.4Hf3.2O17.7) results in a configurational entropy of 1.51 R (R is the gas constant), validating the design as a HEO nanomembrane. The formation of metallic oxides arises from chemical reactions between metallic atoms sputtered from the HEA target and the polymer chains of PVA during the deposition process51,52, facilitating the incorporation of nonmetallic elements, such as oxygen, into the nanomembrane. Furthermore, the abundance of C = C, C–C, and C–O bonds and the organometallic scaffold nanostructures can be attributed to the hydrogel decomposition and the intricate interactions during deposition and exfoliation53. Decomposed residuals were identified through AFM-based nanoscale mapping of Fourier transform infrared spectroscopy (Nano-FTIR, Fig. 2i). The residual C–O and C–O stretching bonds induce stronger absorption at the interphase compared to that of the scaffold (Fig. 2j). The variation in absorption intensity mirrors morphological alternation of the dual phases (Fig. 2k and Supplementary Fig. 10). Based on these observations, we can infer that HEO constitute the scaffold of the nanomembrane, while the interphase predominantly comprises decomposed molecular chains of PVA hydrogel, as schematized in Fig. 2l.

Fig. 2: Characterizations of the dual phases in the HEO scaffold nanomembrane.
figure 2

High-resolution XPS spectra of (a) C 1 s, (b) O 1 s, (c) Co 2p, (d) Ni 2p, (e) Ti 2p, (f) Zr 3 d, (g) Hf 4f , and (h) the corresponding relative atomic percentage of metallic elements with different etching durations. i Nano-FTIR absorption mapping for C–O stretching bonds at a wavelength of 1107 cm-1, (j) point absorption spectra of the dual phases at the corresponding positions and (k) the variation of absorption intensity along the line scan in panel (i). l Schematic illustration for the dual-phase nanostructure. AM-FM mappings of the (m) contact stiffness and (n) loss tangent, and (o) their variations along the line scans in panels (m) and (n). Tk and Tδ denote the correlation lengths of stiffness and loss tangent, respectively. Source data are provided as a Source Data file.

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Moreover, we utilized amplitude modulation-frequency modulation (AM-FM) mapping integrated into AFM to assess the nanoscale variations in mechanical properties (Fig. 2m–o and Supplementary Fig. 11). In Fig. 2m, a substantial contrast in contact stiffness between the scaffold and interphase is evident. While the exact modulus values for the two phases cannot be precisely extracted due to the limitations of the AM-FM technology54, the significant disparity between them is unmistakable. Additionally, Fig. 2n illustrates the corresponding distribution of the loss tangent (the ratio of the loss modulus to the storage modulus). The line scan profiles clearly reveal that the interphase, with lower stiffness, exhibits a higher loss tangent, and vice versa (Fig. 2o). Essentially, the decomposed PVA interphase is compliant yet demonstrates a greater capacity for viscous dissipation during deformation, contrasting sharply with the rigid HEO scaffold. The variations in elastic stiffness and viscous dissipation exhibit similar correlation lengths, and an inverse relationship in spatial distribution. This remarkable heterogeneity in nanoscale viscoelastic properties is absent in conventional single-phase nanomembranes (Supplementary Fig. 12). The in-situ formation of a nanostructured scaffold, interspersed with a soft interphase capable of dissipating energy, reflects the design principles observed in various biological materials (e.g. articular cartilage37, bones38, nacre39), thereby conferring exceptional mechanical attributes to this distinctive nanomembrane.

Mechanical properties and superior fracture resistance

Following the established methods49,55, we conducted AFM-based indentation tests to investigate the mechanical properties of our HEO scaffold nanomembrane, juxtaposed with the HEA and HEO single-phase nanomembrane for comparison (Fig. 3a–c and Supplementary Figs. 13 and 14). Figure 3a presents a comparison of the normalized force-displacement curves, which were analyzed through elastic finite element modeling (FEM) simulations (see the Methods section). Evidently, the HEA and HEO single-phase nanomembranes, devoid of the distinctive dual-phase structure, exhibit common brittle fracture modes observed in 2D materials32,36. This observation is corroborated by their fracture morphologies post-indentation (Fig. 3b), where cracks rapidly propagate from the indentation center to the periphery. In contrast, our HEO scaffold nanomembrane demonstrates a prolonged phase of plastic deformation, followed by a transition to ductile fracture at the indentation center. These phenomena collectively suggest that our nanomembrane exhibits enhanced toughness and greater resistance to fracture. We analyzed the indentation data to extract the mechanical properties of the freestanding nanomembranes (Fig. 3c and Supplementary Fig. 15). The modulus (~3.9 ± 0.2 GPa) and yield strength (~143 ± 38 MPa) of the HEO scaffold nanomembrane are comparatively lower than those of the single-phase nanomembrane, attributed to the incorporation of decomposed molecular chains within the interphase. Additionally, the HEO scaffold nanomembrane achieves a maximum von Mises strain approaching 90% before undergoing mechanical softening, while the toughness, quantified by the area integral under the stress-strain curve, surpasses 300 MJ/m3. This exceptional toughness, particularly notable given the nanoscale thicknesses involved, is a rare attribute in HEO materials34,56. Consequently, the HEO scaffold nanomembrane surpasses a wide range of thin film materials and 2D nanosheets, exhibiting a strength within the metallic range, coupled with superior ductility and toughness akin to polymer films and composites (Fig. 3d, e).

Fig. 3: Superior mechanical performances and hierarchical toughening mechanisms of the HEO scaffold nanomembrane.
figure 3

a Normalized load-depth curves under AFM indentation and (b) the corresponding fracture morphologies of different nanomembranes. c FEM fitting to mechanical properties of the HEO scaffold nanomembrane, and comparison with other materials64,71,72,73,74,75,76,77,78,79,80,81,82,83,84,85,86,87 in terms of (d) yield strength versus ductility, and (e) Young’s modulus versus toughness. BF-STEM image of (f) crack initiation, (g) blunting, and (h) deflection in the ruptured HEO scaffold nanomembrane. i DF-STEM image of the enlarged view for the dashed-rectangle region in panel (h) showing nanobridging, and (j) HRTEM image of the highly deformed interphase. Error bars are the standard deviations. Source data are provided as a Source Data file and in Supplementary Table 1.

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To understand the exceptional fracture resistance, we delved deeper into the fracture process zone in our HEO scaffold nanomembrane. Figure 3f–j display the morphologies of a ruptured nanomembrane, shedding light on its hierarchical toughening mechanisms. As seen in Fig. 3f, the failure initiates with the simultaneous emergence of multiple dispersed cracks. Intriguingly, the ductile interphase mitigates these nanoscale cracks by stretching beyond 300%, effectively impeding crack propagation (Fig. 3g). As crack propagation ensures, the dual-phase nanostructure facilitates continuous crack deflection, with evidence of crack bridging among individual scaffold components, a strong indication of toughening57 (Fig. 3h). Moreover, this bridging phenomenon extends to the nanoscale interphase (Fig. 3i), accompanied by profound deformation and the formation of nanovoids ahead of the crack tip (Fig. 3j). The bridging observed between scaffold components and the distinct interphase underscores the hierarchical toughening mechanisms operating across length scales from a few nanometers to tens of nanometers.

Moreover, we employed finite element simulations to explore the mechanisms of crack initiation and propagation within the HEO scaffold nanomembrane using a thermodynamically consistent phase field model58. The mechanical properties of the materials utilized in these simulations were derived from our experiments (see the “Methods” section). By incorporating a representative TEM image of the HEO scaffold nanomembrane into the simulations to depict the spatial distribution of scaffold and inclusion phases (Fig. 4a), we observed notable distinctions in the simulated stress-strain curves between the dual-phase and single-phase nanomembranes under mode-I cracking (Supplementary Fig. 16a). In the simulations, the deformed dual-phase nanomembrane showcased a gradual softening behavior, contrasting sharply with the abrupt stress drops typically associated with rapid crack propagation in single-phase brittle nanomembranes. The substantial plastic dissipation within the dual-phase nanomembrane contributed to a high fracture energy (Supplementary Fig. 16b). The mismatch in elastic moduli between the scaffold and inclusion phases played a pivotal role in enhancing the material’s toughness through various mechanisms. Primarily, this moduli mismatch created a heterogeneous stress field, triggering non-local plastic yielding and crack initiation (Fig. 4b), mirroring similar dispersed cracking patterns observed in TEM characterizations (Fig. 3f). Additionally, the elastic mismatch reduced the local crack driving force (J-integral) as the crack propagated through a softer inclusion phase59, facilitating crack shielding and deflection (Fig. 4c). Further investigations illustrated that as the elastic mismatch increased, the crack surface roughened, thereby augmenting fracture toughness (Supplementary Fig. 17). Conversely, the absence of crack deflection in the single-phase nanomembrane resulted in poor fracture toughness (Fig. 4d).

Fig. 4: Phase field simulations of the fracture process of the HEO nanomembrane.
figure 4

a Illustration of the cracked model with a dual-phase morphology subject to uniaxial tension along y-axis (uy). b Enlarged view of dispersed crack initiation and soft inclusions bridging during the early stage of crack propagation. c, d Snapshots of the crack propagation for the dual-phase and the single-phase models, respectively. Contours correspond to the normal stress along y-axis (σy).

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Multifunctional nanomembrane and encapsulation performance

Having confirmed the exceptional mechanical properties of our nanomembranes, we proceeded to explore their multifunctional capabilities. A key feature facilitating their wide applicability is their ease of transfer onto various substrates through van der Waals forces60, ensuring effortless handling (Fig. 5a and Supplementary Fig. 18). The complex chemical composition, comprising nano-sized oxides, metals, and decomposed polymer chains, imparts strong adhesion to a range of materials, including metals, polymers, silicon wafers, and glass slides, among others33,61. This robust bonding endures even under demanding conditions—when fully immersed in water and subjected to vigorous mechanical agitation, the nanomembrane steadfastly remains affixed to the substrate (Fig. 5b and Supplementary Movie 2). Furthermore, the pull-off force required to separate the nanomembrane-covered polydimethylsiloxane (PDMS) bulks is approximately four times greater than that for two bare PDMS surfaces, underscoring the excellent bonding strength with the polymer substrate (Fig. 5c and Supplementary Fig. 19). Beyond their strong adhesion, our nanomembranes exhibit superior optical properties. They offer exceptional transparency, enabling clear visualization of underlying flexible circuits, while also reducing the reflectivity of polished silicon wafers, resulting in a darker appearance as illustrated in Fig. 5a. Spectral analyses of transmittance and reflectance within the visible light range reveal an average transmittance of 83.2% and reflectance of 9.2% (Fig. 5d). These optical characteristics closely match those of a standard glass slide, surpassing numerous transparent TFEs used in luminescent or display devices11,15,16. This optical performance can be attributed to the nanoscale thickness, low surface roughness, amorphous natures of the two nano-phases, and the visible-light transparency of metal oxides, typically characterized by a bandgap falling within the range of 3-5eV20.

Fig. 5: Adherent, transparent, stretchable, anti-oxidant HEO scaffold nanomembrane for TFE of flexible devices.
figure 5

a The HEO scaffold nanomembrane on the flexible circuit and Si wafer. b Photographs of the nanomembrane TFE before and after being immersed in water. c Load-displacement curves for pull-off tests of bare and nanomembrane-covered PDMS bulks. d Visible transmittance and reflectance spectra of the nanomembrane on a glass substrate with the bare glass as a reference. The inset shows a photograph of the circuit visible through the nanomembrane. e Photographs and height images of the nanomembrane TFE on the flexible circuit before and after being stretched and bent in water. f Uniaxial tension to crack initiation at a failure strain of εf and (g) fatigue tests of 2000 cycles at 5% strain for the nanomembrane on PDMS. h Photographs and height images of the nanomembrane TFE on Cu film before and after being bent in water. i Height images of the nanomembrane TFE on the Cu film after 24 h of 85 °C/85%RH accelerated life test, and j the corresponding high-resolution XPS spectra of Cu LM2 with and without the nanomembrane TFE. Source data are provided as a Source Data file.

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The combination of these versatile functionalities positions our nanomembrane as a highly promising candidate for a wide array of applications. To showcase its potential, we rigorously evaluated its performance as a TFE for flexible devices. To this end, our nanomembrane was seamlessly integrated into a flexible circuit featuring a patterned Pt film on a PDMS substrate, as shown by the optical and height images captured via confocal laser scanning microscopy (Fig. 5e). When subjected to immersion in water and varying degrees of tensile and bending stress, the nanomembrane TFE exhibited remarkable resilience, maintaining its structural integrity and adhesion without any discernable cracking or failure (Fig. 5e and Supplementary Movie 3). Uniaxial tension tests conducted on the nanomembranes atop PDMS substrates62,63 or water surface64,65 indicated that the macro-crack initiation occurred at approximately 8% strain under monotonic loading while, under cyclic loading, the TFE can endure 5% strain for 2000 cycles without crack initiation or propagation (Fig. 5f and g, and Supplementary Fig. 20). This flexibility under deformation and fatigue remained largely unaffected by substrate contaminants or transfer-induced defects like wrinkling, which is invaluable for their practical utility. The combination of mechanical robustness and the inherent anti-oxidation traits of HEO materials26,27 equips the nanomembrane to offer stable protection for delicate circuits or devices.

To further validate its efficacy, we transferred the nanomembrane onto Cu films and subjected the entire system to immersion and bending it in water with a ~ 5% strain (Supplementary Movie 4). In the absence of the nanomembrane TFE, the Cu film quickly oxidized in the humid environment, evident by the emergence of dark spots depicted in Fig. 5h, whereas the encapsulated Cu film remained largely unaltered. Subsequent accelerated life tests at 85 °C and 85% relative humidity (RH) revealed extensive oxidation of the Cu film, characterized by the formation of reddish copper (I) oxide66 (Fig. 5i). In stark contrast, the nanomembrane TFE significantly mitigated the oxidation, as confirmed by high-resolution XPS and Raman spectra (Fig. 5j and Supplementary Fig. 21). These observations underscore our nanomembrane TFE’s exceptional ability to achieve a low water vapor transmission rate (WVTR), safeguarding flexible devices against failure despite its mere 30-nm thickness. This achievement stands out in comparison to existing TFEs, such as polymer TFE prone to high permeability67, or rigid metal-oxide films lacking deformability17, marking a significant advancement in the realm of ultrathin flexible device development.

Discussion

The heterogeneous HEO scaffold nanostructure of our nanomembranes, born from complex reactions between sputtered chemically complex HEA with the organic substrates of hydrogel49,52, presents a visually captivating mosaic. This structured arrangement, akin to Turing patterns found in biological and chemical systems, emerges through a reaction-diffusion process involving multiple species, stirring thoughts of ordered complexity8,32. Interestingly, the formation of our scaffold nanostructure and HEO component mirrors this spontaneous yet controllable process within our fabrication technique, offering not only a visually arresting pattern but also endowing the inherently brittle HEO materials with exceptional fracture resistance. This approach holds promise for extending its application to other materials, enhancing mechanical properties and amalgamating diverse functionalities.

In essence, we have engineered a fracture-resistant HEO nanomembrane boasting a heterogeneous nanostructure comprising a hard HEO scaffold intertwined with soft decomposed PVA chains. This design exhibits enhanced ductility and toughness, evident in AFM indentation and macroscopic mechanical tests, significantly outperforming their single-phase homogeneous counterparts. Moreover, these nanomembranes exhibit strong and robust adhesion, optical clarity, and inherent anti-oxidative properties. Having demonstrated their ability to meet the stringent requirements of ultrathin flexible encapsulation for mechanical durability and environmental resilience, their potential extends beyond promising applications across various domains due to their exceptional geometrical, physical, and chemical attributes.

Methods

Sample preparation

The HEO scaffold nanomembranes were fabricated via the polymer surface buckling enabled exfoliation technique46,47. Initially, we spin-coated 11.1 wt% PVA solution onto a 10 × 10 cm2 glass substrate for 30 s using different spin-coating speeds, and dried it in an oven at 60 °C for 24 h to obtain the dehydrated hydrogel films. Subsequently, we sputtered a HEA target with a nominal composition of (CoNi)50(TiZrHf)50 (at.%) onto the hydrogel films by magnetron sputtering under a background vacuum of 4E-6 Torr and a working pressure of 2E-2 Torr. The sputtering power of the 3-inch HEA target was controlled at 150 W. For comparison, the HEA target and a HEO target with the composition of the scaffold (Co1.6Ni0.7Ti2.1Zr2.4Hf3.2O17.7) were also sputtered and deposited onto the NaCl substrate instead of the hydrogel substrate, following the same procedure. After that, we cut the deposited films into the desired shape using a blade and immersed the overall system in deionized (DI) water. As the water diffused into the PVA film causing it to swell and surface buckle (the NaCl substrate was dissolved), the freestanding nanomembranes were exfoliated from the substrate. As a result, we could collect the nanomembranes floating in DI water using different substrates, such as TEM grids, silicon wafers, and glass slices, for subsequent characterization, testing, and application.

Structural and compositional characterization

The morphology, atomic structure, and elemental distribution of the prepared nanomembranes were examined using a TEM (JEOL 2100 F) and a STEM (Thermo Scientific Talos F200X) operating at 200 kV. To determine the chemical states of all elements, X-ray photoelectron spectroscopy (Thermo Fisher, ESCALAB Xi + ) with monochromatic Al Kα radiation was employed. An Ar-ion beam operating at 1 keV was used to etch the nanomembrane applying a spot size of 3 × 3 mm² for the depth profile analysis.

The thickness and surface morphology of the nanomembranes on Si wafer were characterized using AFM (MFP-3D Infinity, Oxford Instruments), employing a single crystal diamond AFM probe (ART D160, K-TEK Nanotechnology). Meanwhile, the nanomembranes were investigated in AM-FM viscoelastic mapping mode for nanomechanical properties using an Olympus AC160 probe. The cantilever parameters are calibrated54 as resonant frequencies of 302 and 1717 kHz and spring constants of 33 and 1100 N/m. The correlation lengths were determined through correlation function calculations using Gwyddion software68.

Nano-FTIR spectroscopy of the nanomembranes on Si wafer was conducted using an AFM-based setup (neaSNOM, neaspec) based on probing the laser-induced photothermal expansion69. A PtIr-coated AFM probe (PPP-EFM, Nanosensors) was used for both topography imaging and spectral measurements with a range of 1815–920 cm-1. Nano-FTIR imaging was subsequently performed with a resolution of 1 cm-1 at selected locations according to the peak positions observed in the spectra.

Mechanical characterization and FEM simulation

We investigated the mechanical properties of the nanomembranes by AFM indentation following a well-established protocol49,55 (Supplementary Fig. 13–15). Initially, the freestanding nanomembranes were transferred to Si wafers with cylindrical holes machined by microwell etching. Prior to indentation, AFM (MFP-3D Infinity, Oxford Instruments) was used to image the morphologies of the suspended nanomembranes, and the probe tip was positioned at the hole center. Sequential loading-unloading indentations were then performed with a single-crystal diamond AFM probe (ART D160, K-TEK Nanotechnology, tip radius ~16 nm). The stiffness of the cantilever was calibrated as 6.84 N/m. Force-displacement curves were obtained at an indentation velocity of 200 nm/s, supplemented by morphological observations after each loading to determine the fracture of the nanomembranes. At least 3 independent force-displacement curves were recorded for each sample.

The mechanical properties of the nanomembranes were derived using FEM simulation with commercial software ABAQUS following previous methodologies49. An axisymmetric section was modeled as the suspended nanomembrane, with the thickness and profile determined based on the measured AFM topography. The probe tip was modeled as a frictionless and rigid sphere. Without loss of generality, the Poisson’s ratio was presumed to be 0.3. An isotropic linear-elastic model was employed to extrapolate the elastic modulus by aligning the elastic segment of force-displacement curves with simulation results. After that, we simulated the indentation using an elastoplastic model based on the von Mises yield criterion. The yield stress was defined as the deviation from linear elasticity, and a bilinear isotropic hardening model was used after yielding. The ductility was evaluated as the maximum von Mises strain before mechanical softening, and the toughness was obtained by integrating the area under the stress-strain curves.

Phase field simulation

Simulations of the deformation and fracture behavior of the dual-phase nanomembrane were carried out using COMSOL software. Phase field damage and J2-type plasticity theories were adopted. A quadratic degradation function was used to describe the material softening with increasing damage variable. Uniaxial tension was applied to a 160 × 16 nm2 model with single edge crack (50 × 5 nm2) under plane stress condition. The mesh size is 1 nm. A representative dual-phase morphology extracted from STEM was input in the finite element model. Young’s modulus, Poisson’s ratio and yield stress were determined based on the experiments: Esca = 20 GPa, Einc = 2 GPa, and νsca = νinc = 0.3. The initial yield stress of 200 MPa and the isotropic hardening modulus of 400 MPa were used for the inclusion phase. The critical energy release rate Gc = 1 J/m2 and the crack length scale l = 5 nm were used for the phase field theory. Additionally, the fracture process of the single-phase brittle nanomembrane was simulated using linear elasticity with E = 20 GPa and ν = 0.3. To study the effects of elastic mismatch, we also simulated the fracture process using different modulus ratios (Esca/Einc = 1, 2, 10, 20), while the average modulus remained unchanged.

Thin film encapsulation test

We used a mechanical testing stage (MFS350, Linkam) to evaluate the adhesion between the nanomembrane and PDMS by pull-off tests, and the stretchability of the nanomembranes by macroscopic tensile tests. For the pull-off tests, the nanomembranes with a lateral size of 5 × 5 mm2 were transferred onto PDMS bulks with a dimension of approximately 10 × 10 × 5 mm3. Two of the prepared systems, with their nanomembrane-covered sides facing each other, were subjected to compression with varying loads and durations, followed by separation at a velocity of 20 μm/s. For the uniaxial tensile tests, the nanomembranes were transferred onto PDMS tensile specimens with a gauge of 20 × 6 mm2 and water surface with PDMS clamps64,65. Subsequently, the prepared specimens were stretched at a strain rate of 5 × 10-4 s-1. We also conducted the fatigue test of 2000 cycles at 5 Hz and 5% strain using an electromagnetic dynamic fatigue testing machine. Optical microscope images before and after the mechanical tests were captured using a Leica DM2500 optical microscope with an INFINITY8-8M camera (Teledyne Lumenera).

The optical properties of the nanomembrane were obtained using a UV/Vis/NIR spectrophotometer (Lamda 950, PerkinElmer). The transmittance and reflectance spectra within the visible light range (wavelength ranging from 380 to 800 nm) were measured for the nanomembrane (20 × 10 mm2) on a glass substrate with the bare glass slide as a reference.

For the encapsulation tests, the nanomembranes were transferred onto 3-μm thick electro-deposited Cu films supported by a 25-μm thick polyimide (PI) substrate. The bending strain of the entire system atop an aluminum foil tape of 100 µm thickness can be estimated as ε = t/2R70, where t is the overall thickness and R is the bending radius. The accelerated life tests of the encapsulated samples were then conducted at 85 °C/85%RH conditions for 24 h in a temperature humidity chamber. The Cu films with and without the nanomembrane TFE were observed using a confocal laser scanning microscopy (VK-X100K, Keyence) with a height resolution of 0.005 μm. Subsequently, we measured the XPS spectra and Raman spectra at the corresponding positions to identify the formation of copper oxide. A confocal Raman microscope (inVia Qontor, Renishaw) with an excitation wavelength of 532 nm was used.