Atomic manufacturing process
Figure 1 illustrates the schematic process of obtaining a thin film using PLD as an atomic manufacturing strategy. An overview diagram of the entire procedure is presented in Fig. 1a. The pulsed laser is precisely focused on the target, triggering a plasma plume composed of atomic clusters with varying sizes and energy. These atomic clusters form high-throughput libraries with diverse energy distributions and sizes, as depicted in Fig. 1b. These libraries act as the fundamental building blocks necessary for achieving a wide range of microstructures. The schematic representation in Fig. 1c demonstrates the landing of different atomic clusters on the surface of the film. Supplementary Fig. 1 provides a schematic illustration of the experimental apparatus utilized for atomic manufacturing and sample diagrams of specific structures. The extremely small size of these clusters enables rapid kinetic behavior40,41, facilitating quick structural rearrangements to find more stable configurations and ultimately leading to the formation of densely packed films. The benefits of atomic manufacturing can be elucidated through three key aspects. Firstly, the atomic manufacturing process preserves the diversity of atomic clusters, which is typically lost in conventional rapid solidification methods that disrupt the cluster diversity of the melt. Secondly, the diverse atomic clusters can be more efficiently rearranged upon reaching the film surface, resulting in a more stable and denser microstructure. Thirdly, by adjusting the manufacturing parameters, the atomic cluster libraries can be conveniently modified, allowing for a wide range of conformations and structures to be realized.
Microstructure regulation
Figure 2 illustrates the construction of a spatial framework encompassing materials with varying microstructures. The x-axis represents the characteristic size of the structural units, while the y-axis indicates the degree of order. Traditionally, acquiring such a space necessitates a multitude of materials with distinct compositions or diverse preparation and processing methods. The technique of atomic manufacturing is possible to create diverse microstructures with precise control while keeping the compositions fixed. Consequently, we have successfully generated a series of diverse microstructures, and Fig. 2a-n denotes 14 representative obtained microstructures, which exemplify four primary classes of nanostructures achieved through atomic fabrication within the order/size space. The XRD patterns of different nanostructures are shown in Supplementary Fig. 2. Figure 2a, b present high-resolution transmission electron microscopy (HRTEM) images of the MG, exhibiting a highly disordered structure with atomic-scale characteristic size. This disordered structure is confirmed by the diffuse diffraction ring in the fast Fourier transform (FFT) of the inset, placing it in the red region of the order/size space.
By elevating the pulsed laser intensity during atomic manufacturing, microstructures with enhanced order and larger characteristic sizes can be achieved, represented by the orange regions within the order/size space. Figure 2c, d illustrates a structure resembling spinodal decomposition, characterized by the entanglement of two distinct amorphous nanophases (no obvious bright spots in FFT). This entangled state, referred to as the SG, exhibits a spherical shape with a diameter ranging from 2 to 4 nm, as depicted in Fig. 2c, while in Fig. 2d, the characteristic size is larger, approximately 9 nm, resembling an olive shape. As illustrated in Supplementary Fig. 11, the energy-dispersive spectroscopy (EDS) analysis reveals disparities between the two amorphous phases characterized by variations in the presence or absence of Cu enrichment.
By progressively increasing both the substrate temperature and the pulsed laser energy during atomic manufacturing, a diverse array of DPNs can be obtained. These structures are situated in the green region of the order/size space, corresponding to the junction region of orange and green regions. As depicted in the HRTEM images of Fig. 2e, f, the non-uniformly distributed nanocrystalline grains with dimensions of 1–2 nm precipitate on the amorphous matrix. This distinctive configuration can be considered as an intermediate state in the transition from MG to a DPN. By further increasing the energy and substrate temperature, the nanocrystalline grains exhibit enhanced clarity in their profiles, assuming a characteristic DPN configuration. The morphology of the nanocrystalline grains in the DPN transitions from spherical to elliptical, with characteristic dimensions increasing from 2–3 nm to 8–13 nm. Figure 2g-j collectively presents representative feature structures of DPN, each distinguished by distinct characteristic sizes and orders. This variability signifies the ability to precisely prepare and manipulate DPN structures with desired feature sizes by modifying the atomic manufacturing process.
Furthermore, the characteristic size of nanocrystals increases significantly to tens of nanometers, as depicted in Fig. 2k, l. This is evidenced by the presence of distinct and regularly arranged bright spots in the FFT pattern. At this particular stage, the proportion of the crystalline phase becomes more prominent. This indicates that we can achieve a distinct structure characterized by the prevalence of crystalline phases within the region corresponding to the crystals in the order/size space. Additionally, amorphous phases are observed to be dispersed at the intersection sites of the grain boundaries, as illustrated in Fig. 2m, n. At this point, a well-organized and regularly arranged array, referred to as the CG, is discernible in the FFT analysis. In summary, through meticulous parameter tuning in the atomic manufacturing process, we have successfully achieved diverse microstructures within the order/size space, effectively implementing the concept of elemental simplification. Accordingly, we have identified four distinct structures as potential candidates for a comprehensive examination of their mechanical properties.
Mechanical properties
We conducted nanoindentation and microcompression experiments on micropillars to investigate the mechanical response of four distinct structures (i.e., MG, SG, DPN, and CG) at room temperature. By analyzing the hardness distribution of different nanostructures (Supplementary Fig. 3), a preliminary judgement on their mechanical properties was obtained. The true stress-strain curves, depicted in Fig. 3a, faithfully capture the intrinsic deformation behavior of these characteristic structures. To ensure data consistency and minimize bias, at least three nanopillars for each feature structure and the thickness of the films is more than 1 µm (Supplementary Fig. 4). The engineering stress-strain curves of different nanostructures and the comparison of the morphology before and after deformation are demonstrated in Supplementary Figs. 5–8. One striking observation is the substantial disparity in yield strengths among the four characteristic structures. The yield strength is measured at 2.56 GPa for MG, 3.89 GPa for SG, 3.1 GPa for DPN, and 2.13 GPa for CG, respectively. When substituting the engineering stress-strain curve with the true stress-strain curve, the true strain exceeds 100%, and the strain value corresponding to the engineering strain is chosen. Additionally, all three structures, except MG, exhibit remarkable plastic deformation capabilities exceeding 70%. Notably, SG shows the highest yield strength and notable work-hardening behavior before reaching a strain of 30%. To validate the reliability of work-hardening, unloading was specifically chosen during the microcompression procedure at the hardening stage. Subsequently, reloading was executed, as illustrated in Supplementary Fig. 9, and SEM morphologies were acquired for each stage. To gain deeper insights of the deformation exhibited by samples with distinct feature structures, these samples after microcompression were examined by SEM and the obtained SEM images are presented in Fig. 3b-d. It is evident that the MG undergoes brittle deformation characterized by fracture along the main shear bands, displaying minimal plasticity. In contrast, the SG shows homogeneous deformation without observable main shear bands. The DPN structure shows a certain degree of plastic deformation, retaining prominent main shear bands, and exhibiting intermittent serration behavior in the stress-strain curve, indicating the presence of a few small shear bands. On the other hand, CG exhibits excellent plastic deformation behavior, with numerous pronounced serrations in the stress-strain curve, suggesting the formation and intersection of multiple small shear bands. The high strength exhibited by DPNs stems from their core-shell structure, where cores smaller than 10 nm are enveloped by an amorphous shell spanning several nanometers, resembling a grain boundary. This amorphous shell effectively hampers dislocation motion and grain gliding, successfully counteracting the reverse Hall-Petch effect42. The plasticity of DPNs arises from the hindered propagation of shear bands by the nanocrystals, preventing their expansion and instead promoting the formation of numerous embryonic shear bands, thus sustaining plastic strain. Remarkably, SGs surpass DPNs in both strength and plasticity. A more detailed discussion regarding the underlying deformation mechanisms will be discussed subsequently.
Therefore, through atomic manufacturing, a diverse range of microstructures with tunable mechanical properties have been achieved. Considering the impressive yield strength and plasticity displayed by the samples fabricated using atomic manufacturing, we performed a comprehensive statistical analysis of the mechanical properties of Zr-Cu-Al compositions prepared by alternative approaches. Figure 4a illustrates the normalized yield strength relative to Young's modulus, with data points representing nanopillars with a diameter equal to or less than 500 nm. Most MG systems exhibit remarkable strength but limited plasticity. Conversely, multilayer film structures show some enhancement in plasticity but lack sufficient strength. The heterogeneous structures display a substantial improvement in plasticity while maintaining commendable strength. However, our SG structure surpasses these results in both strength and plasticity, and demonstrates notable toughening and strengthening effects compared to conventional Zr-based MG systems. Figure 4a shows the remarkable mechanical properties achievable through atomic fabrication. Importantly, atomic manufacturing exhibits great potential for tuning properties over wide ranges to create disruptive performance, as exemplified in Fig. 4b. The mechanical properties of the Zr-Cu-Al system, obtained through composition tuning and annealing treatments, primarily occupy the lower left corner of the strength-plasticity diagram. The unique structures obtained through atomic fabrication not only possess outstanding mechanical properties but also offer broad tunability, enabling the creation of diverse structures with excellent properties in multiple aspects. The corresponding yield strength and plasticity modulation ranges for different nanostructures are depicted in Supplementary Fig. 10.
Deformation mechanism
After microcompression experiments, the deformed four distinct structures of MG, SG, DPN, and CG were examined by aberration-corrected scanning transmission electron microscopy (STEM) analysis to investigate the underlying deformation mechanism. Figure 5a-d demonstrates a comparison of the MG before and after deformation. The post-compression HRTEM image (Fig. 5b) revealing the formation of main shear bands leading to fracture, which has also been observed in Fig. 3b. Figure 5c, d depict the schematic representation of the morphology before and after microcompression of the MG. Under applied strain, significant atomic rearrangements occur, leading to the emergence and 45° inclined propagation of primary shear bands (denoted by light blue bands) within the MG. As stress increases, localized shear bands traverse the sample, resulting in crack formation and eventual catastrophic failure43. Thus, constructing unique structures to mitigate shear band instabilities presents an opportunity to circumvent the Achilles' heel of MGs.
At a strain of 75%, the SG sample becomes relatively more homogeneous as the spinodal structure is hardly visible in the HAADF-STEM images (Fig. 5f). To gain deeper insights into underlying mechanism of plastic deformation in SG, the deformation evolution was investigated by HAADF-STEM and EDS imaging under increased strain (as shown in Supplementary Fig. 11). Notably, the compositional fluctuation of Cu elements gradually weakens and eventually becomes homogeneous with the progression of strain. Figure 5g, h illustrate how the dynamic mixing of atoms enhances the mechanical stability of distinct amorphous nanophases during plastic flow, enabling the material to withstand significant strains without experiencing catastrophic fracture. Previous studies have demonstrated that the spinodal decomposition structure of two crystal phases could significantly enhance the strength and plasticity44,45. Here, two inherently robust amorphous nanophases are effectively entangled in SG alloy, resulting in a further enhancement of strength. Unlike crystalline alloys where dislocations serve as primary agents of plasticity, in MG, shear transition zones (STZs) play a crucial role46. While increased strain promotes the activation of numerous STZs, the dynamic atomic mixing between adjacent amorphous nanophases47,48 impedes the formation of STZs into shear band embryos. This kinetic process, coupled with multiple interactions, facilitates the homogenization of amorphous nanophases during deformation49, effectively preventing the extension of shear bands. As a result, the material exhibits improved ductility and thereby averting catastrophic fracture, leading to high strength, even demonstrating a work-hardening effect.
Figure 5i-l delineates the deformation sequence for the DPN nanostructure, where nanocrystalline phases hinder the propagation of the main shear band. As shown in Fig. 5j, l, under the action of external loads, the shear band will inevitably encounter the nanocrystals dispersed in the amorphous matrix during its propagation. Thus, a huge stress field appears at the front of the shear band, resulting in the emergence of numerous radial embryonic shear bands or the segmentation of grains by shear band to dissipate energy (red or cyan dashed square in Fig. 5j). This deformation mode has also been observed in magnesium alloys, exhibiting near-ideal strength42. The ingenious structural design of encapsulating nanoscale crystalline domains within an amorphous matrix acts as a fail-safe against catastrophic fracture. The amorphous shell surrounding the crystalline grains acts as a formidable barrier, impeding harmful dislocation dynamics and grain movement, and thus forestalling the reverse Hall-Patch effect that compromises strength in conventional nanocrystalline materials. Even if embryonic shear bands emerge within the amorphous substrate, propagation is obstructed by the embedded crystallites. This restrictive effect of nanocrystals on shear band proliferation is two-fold. Firstly, the interfaces between the crystalline and amorphous phases facilitate the formation of multiple initial shear bands, promoting strain delocalization. Secondly, mature shear fronts impinge on particle boundaries, arresting further propagation. By hindering the runaway growth of isolated shear bands, catastrophic failure is averted.
The deformation mechanism of the CG structure is comprehensively explained in Fig. 5m-p. Compared to the initial state (Fig. 5m), the compressed crystalline phase undergoes drastic deformation, as demonstrated in Fig. 5n. However, the presence of the amorphous phase buffers the drastic deformation. The absence of partial dislocation in fully deformed nanograins suggests dislocation generation and subsequent annihilation at the interface between the crystalline and amorphous phases, as delineated by the white dashed line in Fig. 5n. Nanoscale amorphous regions confined within shear bands utilize their inherent plasticity, enabling ductile behavior at crystalline boundaries. Dislocation dissociation at amorphous-grain interfaces prevents stress concentration from dislocation pile-ups, thereby improving plasticity50. Compelling evidence supports the notation that nano-sized amorphous regions effectively suppress dislocation accumulation, creating an environment conducive to the plastic deformation of the CG51,52. The nanoscale amorphous domains serve as effective sinks for dislocations. As dislocations traverse the composite microstructure, the interspersed amorphous regions adeptly capture these defects, curtailing their unconstrained multiplication. Consequently, the material avoids the runaway dislocation accumulation that hastens mechanical failure.
The deformation study of the four distinct structures, MG, SG, DPN, and CG, indicates the existence of a diverse range of deformation mechanisms in the unique nanostructures achieved through atomic manufacturing. This provides significant opportunities for regulating and optimizing material properties. Overall, atomic-scale manufacturing empowers precise manipulation of nano-architectures to achieve customized mechanical performance. By designing amorphous, crystalline, and interfacial constituents in consummate harmony, nanostructures can exhibit various deformation mechanisms, breaking the traditional trade-off dilemma between strength and ductility while maximizing specific properties.
In summary, by harnessing the capabilities of atomic manufacturing as a bottom-up approach, precise manipulation of atom clusters allows for the creation of a wide range of disruptive nanostructures in the typical Zr50Cu40Al10 composition. Among these structures, 14 representative characteristic structures with varying degrees of order have been identified, falling into four major categories: MGs, SGs, DPNs, and CGs. The fabrication processes can be precisely controlled by parameters such as pulsed laser energy and substrate temperature. These unique nanostructures exhibit significantly enhanced mechanical properties, particularly the newly synthesized SG structure, which achieves an unprecedented strength of 3.89 GPa in Zr-based alloys simultaneously exhibiting remarkable plasticity exceeding 75%. This groundbreaking achievement disrupts the traditional trade-off relationship between strength and plasticity. Additionally, the different nanostructured morphologies realized by atomic manufacturing demonstrate a variety of operative deformation mechanisms and provide a diverse range of opportunities to create various unique nanostructures. Particularly important, the refinement of the spinodal structures and dynamic atomic intermixing effectively mitigate interfacial failure caused by stress concentration, and provide a strategy for simultaneous enhancements in both strength and plasticity. This work demonstrates the potential of atomic manufacturing as a methodology for producing controlled and diverse nanostructures, thereby optimizing the mechanical properties of specific composition systems. This advancement not only broadens the range of applications for MGs but also offers novel insights for designing materials with exceptional properties. This expands the property space accessible to materials.