Heterostructured multi-principal element alloys prepared by laser-based techniques

Laser surface treatment

LST is widely used in materials science to prepare heterostructured surfaces on top of substrate materials. This technique offers high repeatability and controllability, enabling precise tailoring of processed parts to meet specific requirements. As shown in Figure 1, typical LST techniques are illustrated in three processing configurations according to the difference of raw materials and processing ways. By using laser energy as a heating source or shock wave, LST can control the area and depth of treatment with high accuracy, ensuring consistent quality and performance across various applications^[62-64]. The technique is also beneficial in applications where traditional treatment methods may fall short, as it minimizes thermal distortion and allows for localized treatment without affecting the bulk material properties.

Figure 1. Processing configurations of typical LST techniques including laser surface remelting, laser glazing, laser cladding, laser surface alloying and laser shock peening.

Laser surface remelting (LSR) uses a high-energy laser beam to quickly melt and solidify the surface of a material^[65]. By adjusting the laser’s scanning speed and power, it is possible to control the depth and extent of melting. Not only can LSR create surface layer structures, but also it can create surface gradient structures. A smooth transition in microstructure from the surface layer to the matrix can be achieved by LSR, resulting in excellent overall properties^[66].

Laser glazing is similar to LSR but focuses on creating a thin, uniform layer on the surface, typically between 0.5 and 100 µm thick. This process involves very rapid cooling, often faster than 100,000 degrees per second, which helps form a glass-like or ultrafine grain structure^[50]. Laser glazing can be engineered to produce thermal barrier coatings that are impermeable and resistant to surface corrosion and wear yet maintain porosity in the bulk. In contrast to the superficial layer created by laser glazing, LSR usually produces a deeper molten layer with a typical thickness between 100 and 500 µm, enhancing the ability to form more complex microstructure at localized areas. For instance, this deeper treatment is especially useful for developing gradient materials, where a change in microstructure and properties as desired will be precisely controlled along the building direction of processed materials.

Laser cladding (LC) is another cutting-edge technology used to modify the surface properties of materials. This technique involves depositing cladding material onto a substrate and using a high-energy laser beam to melt this additional material into a thin layer on top of substrate materials^[67]. This process forms a metallurgically bonded layer. Metallic powders, typically delivered through coaxial or lateral nozzles, are commonly used in this method. The interaction between the metal powder and the laser beam generates a melt pool, which solidifies into a durable metal track as the substrate moves^[68].

Laser surface alloying, while resembling LC, primarily modifies the chemical composition of the substrate's surface and its immediate vicinity, rather than adding a distinct layer on top. This process is more complex and requires precise control over the laser parameters and the addition of alloying materials to achieve uniformity and optimal performance of the alloyed layer.

The alloying process includes melting, mixing, and rapid solidification^[50]. The coatings that contain the alloying elements might be pre-applied using techniques such as electroplating or vapor deposition. During laser treatment, these pre-applied coatings are melted together with the substrate surface, allowing the alloying elements to integrate directly into the melt pool. In contrast, co-deposition refers to the simultaneous feeding of alloy powders or mixtures directly into the melt pool during laser processing. This method ensures that the alloying materials are introduced and mixed in real-time with the molten substrate, potentially offering more immediate and dynamic control over the composition and properties of the resulting alloyed layer. To achieve a uniform microstructure, both pre-deposition and co-deposition methods require overlapping the melting tracks by 20%-30%. This overlap ensures consistent properties throughout the laser-alloyed zone, which differs significantly from the substrate composition^[50].

Laser shock peening (LSP), on the other hand, operates distinctly from other LST techniques by focusing on the mechanical rather than thermal effects. Unlike methods that rely on surface temperatures reaching T_m (the melting point), LSP is effective in avoiding the negative impacts of overmuch heating. During LSP, a plasma shock wave is generated through a high-power laser pulse^[69]. When the material is irradiated by a high-power laser pulse, moderate-temperature and high-pressure plasma is instantly generated. The rapid expansion of this plasma creates a potent shock wave that propagates into the material, inducing high residual compressive stress and plastic deformation on the surface. Consequently, a gradient hardening layer is formed. Without the thermal damage and phase changes associated with higher temperature^[69], this process creates near-surface compressive residual stresses and work-hardening states, which significantly improve resistance to crack initiation and propagation.

LSP is particularly well-suited for the preparation of heterostructured materials with a gradient distribution of hardness and strength, which is highly effective in enhancing the fatigue life of metallic materials^[67]. The high-strain-rate deformation caused by the laser shock pressure (typically in the range of 10⁵ to 10⁷ s^-1) results in greater magnitude and depth of compressive residual stresses and higher strain-hardening ratios compared to other plastic deformation processes such as shot peening and deep rolling^[70]. This makes LSP more effective and widely applied in industry.

Laser additive manufacturing

AM is a transformative approach to creating heterostructured materials by building parts layer by layer according to digital models. This technique allows for intricate designs and precise control over material composition and microstructure, making it ideal for fabricating complex, high-performance MPEA components^[71-73].

Advancements in control systems have significantly enhanced the economic efficiency of producing parts via layer-by-layer AM techniques. LAM exemplifies this progress. As depicted in Figure 2, it is a technique that fabricates three-dimensional objects by sequentially adding material^[74]. Selective laser melting (SLM) stands out as a prominent method within LAM, further showcasing the capabilities of this technology in precise and efficient manufacturing. First, the 3D model is decomposed into a series of two-dimensional layers that mimic the final product shape. Then, a laser beam scans these two-dimensional cross-sections, melting metal powder and forming particle bonds, layer by layer, to construct the final three-dimensional part^[53,75]. The speed of SLM is generally slower than some other AM techniques due to the need for precise layer-by-layer scanning and residual stresses and warping may be introduced during the process, requiring post-processing sometimes.

Figure 2. Effective LAM techniques for creating heterostructured materials including selective laser melting/sintering and laser direct deposition.

Direct metal laser sintering (DMLS) is another key AM technique that utilizes a high-wattage laser to fuse powdered metals into solid structures. Unlike SLM, DMLS sinters the metal powder without completely melting it, which helps build objects layer by layer with fine detail and excellent surface finish^[76]. This method is particularly useful for creating porous metal components and complex geometries. Additionally, DMLS can be adapted to sinter other materials such as ceramics and polymers, a process known as selective laser sintering (SLS). A major advantage of DMLS and SLS is their ability to produce parts with minimal residual stresses and internal defects due to partial melting, which are common drawbacks in traditionally manufactured components^[67]. Nevertheless, when compared with SLM, the application scope of DMLS is relatively restricted. The reason is that the material density and mechanical properties obtained through DMLS are generally somewhat inferior.

Other common LAM methods include laser direct deposition (LDD). It is a type of Directed Energy Deposition, which covers a range of AM technologies that use energy sources such as lasers, electron beams, or plasma arcs. LDD involves directly feeding metal powder or wire material into the laser focal area, melting it, and depositing it onto the substrate to form layered structures^[77,78]. It often uses higher laser speeds to maintain a continuous flow of material and rapid deposition, which can enhance build speed but may compromise precision and surface quality. This method is suitable for repairing and manufacturing complex parts^[79]. By precisely controlling laser power, scanning speed, and powder supply parameters, different material compositions and performance variations can be achieved within the same component. For example, in SLM and LDD processes, gradient structures can be created by altering the type and proportion of powder material, resulting in diverse mechanical properties and corrosion resistance in different regions^[74,80]. This is particularly important for manufacturing complex parts with high performance and multifunctionality.

When comparing these techniques, SLM typically involves slower laser speeds to ensure precise melting and solidification, leading to high accuracy but longer build times. This method often introduces residual stresses and warping, requiring post-processing. In contrast, DMLS/SLS uses moderate laser speeds, balancing between precision and build time, as sintering requires controlled heating but not complete melting. This results in fine details and excellent surface finishes with minimal residual stresses. LDD utilizes higher laser speeds to maintain a continuous flow of material and rapid deposition, enhancing build speed but potentially compromising precision and surface quality.

The AM technique is characterized by unique forming modes, such as melt pool connection, layer-by-layer deposition, high cooling rates, and nonequilibrium solidification. These processes often result in heterogeneous structures with distinctive features, including melt pools, columnar grains, lamellar structures, core-shell structures, phase heterogeneity, and nano-precipitation^{[24,52,81,82]}. These microstructural characteristics contribute significantly to the enhanced mechanical properties and performance of the final components. While each laser-based AM technique offers distinct advantages, they also confront some specific challenges. These include issues related to residual stresses, porosity, and the need for post-processing to achieve desired surface finishes and dimensional accuracies. The choice of technique depends on the desired balance between precision, speed, material properties, and cost, making it essential to tailor the method to specific application requirements.

Gradient structure

A grain size gradient is one of the typical characteristic features of gradient materials. As a traditional mechanical method, severe surface plastic deformation is commonly employed to refine surface grains, resulting in an increasing grain size gradient from the surface to the center. LSP operates on a similar principle. The study by Fu et al. confirmed the effectiveness of LSP in introducing gradient structures to enhance the mechanical properties of MPEAs^[86]. As illustrated in Figure 4A and B, the LSP-treated CrFeCoNiMn_0.75Cu_0.25 MPEA develops a hardened surface layer of several hundred micrometers thick. Notably, the sample that goes through four cycles of LSP treatment exhibits significant surface grain refinement and marked improvements in mechanical properties.

Figure 4. Gradient structures created by LST techniques: (A and B) Cross-sectional EBSD images of CrFeCoNiMn_0.75Cu_0.25 MPEA processed by 1 and 4 cycles of LSP, respectively^[86]. (C and D) Grain size distribution of as-fully recrystallized (CrCoNi)₉₂Si₄B₄ MPEA after and before LSR^[87]. (E1 and E2) grain distribution of inverse gradient LSR-CoCrFeMnNi, more details are given in (F) and (G) for clear view of local microstructures containing heat-affected zone (HAZ) and fusion zone (FZ)^[89].

Another widely used laser treatment method is LSR. Unlike LSP, which typically produces a thin surface layer composed of fine crystals, LSR generates a thick molten layer due to the greater penetration depth of the laser. The resulting gradient structure is highly dependent on the specific laser parameters employed. Yuan et al. demonstrated that LSR can produce a hardened and corrosion-resistant surface layer in a fully recrystallized (CrCoNi)₉₂Si₄B₄ MPEA^[87], as illustrated in Figure 4C. The microstructure observed includes characteristically oriented columnar grains, perpendicular to the molten pool boundary, with an average grain size of approximately 25 µm - significantly larger than that of the original base material [Figure 4D].

By combining laser processing with deformation, researchers have linked the resulting gradient structure to recovery and recrystallization, resulting in distinct structures. Zhang et al. achieved inverse gradient grain CoCrFeMnNi MPEAs, demonstrating exceptional strength-ductility combinations through a cold rolling-laser surface heat treatment method, which shows a grain size progression from the tough core to the soft surface^[88]. Subsequently, Shen et al. employed LSR on cold-rolled CoCrFeMnNi MPEAs using single and double-sided laser passes^[89]. As shown in Figure 4E-G, the single-sided laser-treated MPEA exhibits a heterogeneous gradient structure as follows: (i) coarse columnar grains in the fusion zone; (ii) coarser grains in the heat-affected zone adjacent to the fusion zone, compared to the base material; (iii) a mix of smaller revertant and fine recrystallized grains in the HAZ near the base material; and (iv) distinct highly deformed pancake-like grains.

Building on the concept of inverse gradient, a thicker inverse gradient layer can be achieved through LDD, also known as laser metal deposition. This process involves laser treatment and powder melting deposition. LDD is particularly effective in manufacturing gradient structures because it allows for the controlled deposition of the same or different materials within a single component.

In the study by Gu et al., LDD was applied to heavily cold-rolled CoCrFeMnNi sheets, resulting in a material with a combination of different grain structures across its depth^[90]. This includes four distinct regions: a coarse-grained outer layer, a transitional layer with columnar grains, an inner fine-grained layer, and a partially recrystallized core featuring nanosized grains, as depicted in Figure 5A-C. This strategic manipulation of thermal gradients during LDD leads to a specific reverse gradient structure.

Figure 5. Gradient structures created by LAM techniques: Inverse pole figure (IPF) maps of inverse-gradient structured LDD-CoCrFeMnNi MPEA at (A) region including both the center and surface regions; (B) magnified region near the surface; and (C) magnified region near the center^[90]. (D) Backscattered Electron (BSE) image with a grain orientation map obtained by EBSD of chemical composition gradient TiZrNbTa^[91].

Thanks to the precise control offered by the laser control system, LDD can achieve a chemical composition gradient by varying the composition of the deposited material. Dobbelstein et al. produced compositional gradient refractory MPEAs for screening purposes by in-situ alloying of elemental powder mixtures^[91]. By gradually replacing Zr powder with Nb powder, they obtained a compositional gradient from Ti₂₅Zr₅₀Nb₀Ta₂₅ to Ti₂₅Zr₀Nb₅₀Ta₂₅. As shown in Figure 5D, the mean grain size of the body-centered cubic matrix increases from approximately 2 µm for Ti₂₅Zr₅₀Nb₀Ta₂₅ near the Mo-substrate to about 60 µm for Ti₂₅Zr₀Nb₅₀Ta₂₅ at the tip of the wall structure, which is more influenced by the chemical segregation of Ta in Zr-rich regions than by the cooling rate.

Dual-phase or multi-phase structure

Different from the gradient structures with gradual changes in the composition or/and phases along one specific direction, dual-phase or multi-phase structure is another most common type of heretostructure. This structure not only allows for a more significant performance comparison between different phases but also enhances the toughness, strength, and corrosion resistance of the material through their interactions.

Building on the bidirectional diffusion behavior of Fe, Co, Cr, Ni, and W elements, along with the principles of lattice distortion, Guan et al. successfully prepared FeCoCrNi-W coatings featuring an in situ Laves phase and face-centered cubic dual-phase structure through LC^[92]. The composition of the Laves phase decreases from 39.15% at the surface to 12.67% at the substrate, creating a heterogeneous structure that transitions from dual-phase to eutectic and hypoeutectic [Figure 6A-D]. The Laves phase contributes high hardness and deformation resistance, while the face-centered cubic phase offers excellent plastic deformation capability. This orderly arrangement of hard and soft phases enhances the overall performance of the coating.

Figure 6. Dual-phase or multi-phase structures created by LST and LAM techniques: (A) The solidification process of LC-(FeCoCrNi + W) gradient composite coating and the microstructures formed in the corresponding coating layer; (B-D) Cross-sections showing the spatial distribution of face-centered cubic and laves phases of different depth of LC-(FeCoCrNi + W)^[92]. (E-L) TEM characterization of the region at a depth of 40-60 µm (E and F); 30-40 µm (G and H); 20-30 µm (I and J) and less than 10 µm (K and L) below LSR treated TiZrHfTaNb surface^[93].

For original single-phase MPEAs, laser processing can still facilitate phase transformation, introducing a multi-phase structure into the system and enhancing the synergy of mechanical properties. For instance, Luo et al. prepared gradient nanostructured (GNS) layers on TiZrHfTaNb MPEAs through LSR^[93]. This process allows the original coarse-grained single-phase body-centered cubic TiZrHfTaNb refractory MPEA to gradually decompose into a TiNb-rich body-centered cubic phase, a TaNb-rich body-centered cubic phase, a ZrHf-rich hexagonal close-packed phase, and a TiZrHf-rich face-centered cubic phase. These phases exhibit a gradient distribution in grain size along the depth direction during the grain refinement process, as depicted in Figure 6E-L. The introduction of decomposed multi-phase and gradient nanostructures greatly improves the wear resistance of the alloy. The synergistic effect of these heterogeneous microstructures, spanning from nanometer to sub-millimeter scales, contributes to the alloy's superior properties.

Core-shell structure and skeleton architecture

When exploring the advantages of biphasic or polyphase structures, it is notable that core-shell structure materials have gradually become popular with the further development of material design. The core-shell structure owns more dense textures and provides a more flexible way to adjust performance by combining the excellent properties of different components at micro levels.

Zhang et al. fabricated an AlCoCrFeNi_2.1 eutectic MPEA featuring a unique composite core-shell structure through SLM^[94]. This alloy exhibits a distinctive fish-scale microstructure, comprising ultrafine honeycomb and columnar sub-grains as shown in Figure 7A-H. It presents a dual-phase structure with face-centered cubic and B2 phases, where the B2 phase forms a unique shell configuration around the face-centered cubic phase cell. This specialized microstructure endows the alloy with high strength and good ductility, achieving impressive mechanical properties, including a tensile strength of 1,159.4 MPa and a plastic elongation of 29.0%.

Figure 7. Core-shell structures created by LAM techniques: (A and B) Microstructure perpendicular to building direction (BD) of the SLM-AlCoCrFeNi_2.1; (C and D) Microstructure parallel to BD; (E and F) The micrograph of core-shell structure; (G) The micrograph of lamellar structure; (H) The micrograph of nanoprecipitate and bridge^[94]. (I) pseudo-3D microstructure of SLM-Al_0.5CrCoFeNi; (J) Quantitative compositional analysis by atom probe tomography (APT) shows the distribution of different elements in the face-centered cubic cells and B2 phase on the boundaries; (K) IPF map showing the near-uniform crystallographic orientation in different grains; (L) the corresponding phase map showing the consistent distribution of the B2 (body-centered cubic) phase on the face-centered cubic cell boundaries; (M) TEM images show the morphology of the B2 phase on the boundaries, with the magnified images from locations 1, 2, and 3 showing the nanometer-scale bridges (nano-bridges), where high-density dislocations are observed^[95].

Similarly, Kumar et al. employed SLM to fabricate Al_0.5CrCoFeNi MPEAs, which display a face-centered cubic +B2 nano-bridged honeycomb structure^[95] [Figure 7I-M]. The cell boundaries are formed by the B2 phase, interconnected by nano-bridges with a high dislocation density. The B2 phase contributes strength through solid-solution strengthening and high dislocation density, while the nano-bridges of dislocations connecting the face-centered cubic cells provide pathways for dislocation movement away from the crack tip.

High-density dislocations are commonly observed in metals produced through AM, primarily believed to arise from the significant thermal stresses generated during the solidification process. Based on that, the dislocation-precipitate skeleton structure enhances the mechanical properties of the material by stabilizing the dislocations and controlling the distribution of the precipitated phases. These innovative structures not only optimize material properties at the microscopic level but also achieve higher strength and toughness by fine-tuning the phase interfaces and defect distributions.

Using SLM technology, Mu et al. successfully combined a high-density dislocation structure with a high-volume fraction of malleable nanoprecipitates in the Fe_28.0Co_29.5Ni_27.5Al_8.5Ti_6.5 MPEA^[96]. As depicted in Figure 8, this material presents a novel dislocation-precipitate skeleton structure embedded with high-density malleable nanoprecipitates, exhibiting an impressive tensile strength of approximately 1.8 GPa and a maximum elongation of about 16%. The ultra-high strength primarily results from the synergistic strengthening effects of dislocations and precipitates, while the notable ductility is attributed to the evolution of multiple stacking fault structures.

Figure 8. Skeleton architecture created by LAM techniques: (A) Solidification process of SLM-Fe_28.0Co_29.5Ni_27.5Al_8.5Ti_6.5; (B) Bright field (BF) scanning transmission electron microscopy (STEM) image of prepared showing the high-density dislocations network with the selected area electron diffraction (SAED) pattern from (011) zone axis showing the corresponding microstructure; (C) BF STEM image showing the sub-grain architectures consist of the face-centered cubic disordered multicomponent matrix (DOMCM) phase and L1₂ ordered multicomponent nanoprecipitate (OMCNP) phase, with SAED pattern from (001) zone axis showing the corresponding microstructure; (D) BF STEM image showing the sub-grains composed of face-centered cubic , L1₂ and a small amount of L2₁ phase; (E) The enlarged view of the local sub-grain feature area. (F) The high-resolution high-angle annular dark-field imaging (HADDF) -STEM image of L2₁; (G) The dislocation-precipitate skeleton architecture^[96].

Moreover, the dislocation skeleton slows the movement of dislocations during deformation without completely obstructing their motion, thus maintaining structural stability and preventing premature failure due to boundary stress concentrations. This balance between strength and ductility illustrates the potential of tailored microstructures in enhancing the performance of MPEAs. This greatly benefits the overall performance of the material, though specific material systems must be considered.

Lamellar structure

Another common structure observed in AM is the lamellar structure, which consists of alternating layers of different phases, providing unique mechanical advantages.

For instance, Ren et al. prepared a dual-phase AlCoCrFeNi_2.1 MPEA using SLM, achieving remarkable mechanical properties with a yield strength of up to 1.3 GPa and a uniform elongation of 14%^[35]. This exceptional performance is attributed to its unique layered nanoscale microstructure, which features alternating nanoscale face-centered cubic and body-centered cubic phases, along with semi-coherent interfaces and chemical segregation of components at the nanometer scale [Figure 9A-D]. The high strength results from the nanolayered structure and elevated print-induced fault density, while significant plastic deformation is facilitated by the high strain-hardening capacity of the body-centered cubic phase. Additionally, nanoscale chemical segregation, varying orientations in adjacent regions, and continuous strain gradients further enhance the strain hardening of the body-centered cubic phase.

Figure 9. Lamellar structures created by LAM techniques: (A) IPF map of SLM-AlCoCrFeNi_2.1, showing a magnified local region; (B) Secondary electron micrograph of the nanolamellar structure; (C) Bright-field TEM image of the body-centered cubic and face-centered cubic nanolamellae with PED patterns; (D) HAADF-STEM image showing the modulated nanostructures within body-centered cubic lamellae^[35]. (E) BSE images of LDD-AlCoCrFeNi_2.1; (F) The phases distribution; (G) TEM image of sample; (H) The corresponding SAED pattern taken from the A zone of (G); (I) The corresponding SAED pattern taken from the B zone of (G); (J) The elemental distributions of the frames in (G)^[97]. (K) SEM image of LDD-CoCrNi/(CoCrNi)₈₆Al₇Ti₇ on building direction-sectional direction (BD-SD) plane; (L) EPMA mapping; (M and N) Typical BF-TEM images of CoCrNi layer and (CoCrNi)₈₆Al₇Ti₇ (denoted as Al₇Ti₇) layer in the sample, respectively^[98].

Another example is that Huang et al. utilized LDD to fabricate the AlCoCrFeNi_2.1 [Figure 9E-J] MPEA^[97]. Similarly, the MPEA produced by LDD exhibits a face-centered cubic/body-centered cubic dual-phase eutectic microstructure. This fine and uniform eutectic structure significantly improves the strain hardening and dislocation accumulation capacity of the EMPEA, thereby enhancing its mechanical properties. The LDD process capitalizes on the high cooling rates and eutectic characteristics to form a refined microstructure, resulting in excellent mechanical performance.

In addition, the precise control offered by LAM allows different materials to be combined for achieving artificial lamellar structures. For example, Sun et al. used LDD to create a multilayer heterostructured MPEA alloy. This alloy consists of alternating entropy-intermediate alloys of CoCrNi and (CoCrNi)₈₆Al₇Ti₇^[98][Figure 9K-N]. After appropriate heat treatment, the heterostructure alloy achieves an ultimate tensile strength of over 1 GPa and an elongation at break of 50.6%. These results are much higher than those reported for medium or high entropy alloys in other studies. The improved performance is mainly due to the formation of a nanoscale L1₂ phase in the CoCrNi layer during heat treatment. Additionally, the gradient transition of the eutectoid in the (CoCrNi)₈₆Al₇Ti₇ layer helps prevent the formation of brittle nanoscale lamellar structures. The composition gradient and low strain gradient at the heterogeneous interface work together to enhance both strength and toughness.

Nano precipitates

In the heterostructures prepared using laser-based technology, precipitates play a crucial role especially at nanoscale. These nano precipitates can be uniformly distributed within the matrix by precisely controlling laser parameters and material compositions, significantly enhancing the mechanical properties and corrosion resistance of the materials. The formation mechanism of nano precipitates involves phase separation and the precipitation of solid solutions, resulting in superior properties especially in the enhancement of strength. For instance, Mu et al. utilized SLM to prepare Fe₂₈Co₂₉Ni_27.5Al_8.5Ti_6.5 MPEA samples^[99]. Through short aging heat treatment, they successfully retained a high-density dislocation structure while simultaneously inducing the formation of duplex ordered nanoscale precipitates (L1₂ and L2₁), as shown in Figure 10A-D. These nanoprecipitates are uniformly distributed within the subgrain boundaries and grains, significantly improving the strength and ductility. The combination of these nanoprecipitates and the in-situ dislocation network promotes the strain hardening mechanism, resulting in an impressive ultra-high strength of 1,430 MPa and an elongation of nearly 16%.

Figure 10. Heterostructures with nano precipitates created by LAM and LST techniques: (A) SEM image showing the orientation arranged nanostructure complex parallel to the BD of SLM-Fe₂₈Co₂₉Ni_27.5Al_8.5Ti_6.5; (B) SEM image showing the multiple nanoprecipitates manipulation; (C) TEM image showing the L1₂, L2₁, and face-centered cubic matrix phases; (D) Schematic illustration of the orientation arranged nanostructure complex composed of multiple nanoprecipitates manipulation and in-situ dislocation network^[99]. (E) BF TEM images of the LSR & aged AlCoCrFeNi_2.1 MPEAs; Dark field (DF) TEM images of the face-centered cubic phase (E1) and B2 phase (E2) from the superlattice spot from their selected area diffraction patters (SADPs) shown in the inset. (F) Schematic diagrams of the surface-strengthening process^[100].

In another study, Miao et al. enhanced the surface properties of AlCoCrFeNi_2.1 MPEA through LSR and aging heat treatment^[100]. After LSR, the surface microhardness of the alloy increases by 59%, while the average friction coefficient and wear rate decrease by 26% and 68%, respectively. This improvement is attributed to the reduction in the lamellar spacing of the face-centered cubic and body-centered cubic (B2) phases and the formation of consistent nanoprecipitates within both phases as depicted in Figure 10E.

The remelting process significantly refines the alloy's microstructure, promoting the formation of nanoprecipitates in both the face-centered cubic and body-centered cubic phases [Figure 10F]. Subsequent aging treatment at 600 °C enhances the quantity and distribution of these precipitates, further enhancing the alloy's hardness and wear resistance. Although rapid cooling during laser remelting initially inhibits precipitate formation, the subsequent heat treatment allows for an even distribution and increased strength of the precipitates. The presence of these nanoprecipitates increases the phase interfaces, enhancing the mechanical properties of the alloy, particularly during plastic deformation.

These findings highlight the importance of nanoprecipitates in improving the properties of MPEAs and offer valuable insights for optimizing alloy microstructures at nanoscale.

Hierarchically heterogeneous structure

Building on the previously discussed concepts, it is evident that the rapid cooling characteristics of laser processing technology, combined with the inherent complexity of MPEA systems, often result in a hierarchically heterogeneous microstructure. This type of structure incorporates multiple phases and forms, enhancing material properties. Actually, some of the aforementioned cases belong to the category of hierarchically heterogeneous structure, such as dual-phase or multi-phase with gradient distribution of composition or grain sizes, along with nanoparticipates embedded in lamellar structures. Herein, more typical hierarchically heterogeneous structures will be introduced to broaden our horizons.

For example, Luo et al. investigated a TiZrHfTaNb_0.5 MPEA and highlighted this phenomenon^[101]. The as-cast MPEA exhibits a single body-centered cubic solid solution phase with an average grain size of approximately 115 µm. However, after LSR treatment, a GNS layer about 100 µm thick formed on the surface, showcasing a gradual refinement of grain size with decreasing depth, as illustrated in Figure 11A-J. Within this GNS layer, remarkable microstructural evolution occurred, including the formation of ultrafine grains, precipitated nanostructures, and a unique crystalline-amorphous heterostructure at the topmost layer.

Figure 11. Hierarchically heterogeneous structures created by LAM and LST techniques: (A) SEM image showing the morphology along the depth direction of the LSR-TiZrHfTaNb_0.5; (B) Variations of grain size and microhardness along the depth away from the topmost surface; TEM characterization of crystalline-amorphous nanostructured surface layer (C and D) and at a depth of 10-20 µm (E and F); 20-40 µm (G and H) and 40-60 µm (I and J) below the surface^[101]. (K) IPF map of building plane of the SLM-Fe₆₀Co₁₅Ni₁₅Cr₁₀; (L) The pseudo-3D EBSD KAM maps; (M) SEM. (N) STEM image of solidification cells; (O) XRD pattern with inset showing EBSD phase maps on the plane normal to the scan direction; (P) 3D reconstructed element distributions near the cell boundaries by APT analysis; (Q) 1D concentration profiles of the 3D element distribution in the APT maps^[102].

The development of this complex microstructural architecture is attributed to phase decomposition-mediated gradient refinement mechanisms. The high-energy LSR process melts the outermost surface, creating an amorphous state. Subsequently, rapid and gradient cooling rate during solidification leads to partial crystallization, resulting in the distinctive crystalline-amorphous nanostructure.

Further analysis revealed that compositional segregation and phase transformations significantly influence the evolution of the GNS layer. The initial body-centered cubic solid solution matrix underwent various phase decompositions, resulting in TiTaNb-rich body-centered cubic, ZrHf-rich body-centered cubic, ZrHf-rich face-centered cubic, and minor hexagonal close-packed phases, with their distribution and grain sizes varying along the depth. These compositional and structural changes are closely related to the gradient thermal conditions - temperature and cooling rate - experienced during the LSR process.

Additionally, Park et al. demonstrated the excellent mechanical properties of an iron-based medium-entropy alloy Fe₆₀Co₁₅Ni₁₅Cr₁₀ fabricated using SLM^[102]. The SLM process introduces spatial and compositional heterogeneity into the alloy, characterized by anisotropic grain morphology, a honeycomb submicroscopic cell structure, and compositional segregation at cell boundaries [Figure 11K-Q]. The size and morphology of grains differ significantly in various directions, while the honeycomb cell structure contributes to a unique microstructure that enhances both strength and toughness.

Significant compositional segregation, particularly at cell boundaries, results in variations in chemical composition across different regions. This segregation reduces the stability of the face-centered cubic phase and promotes strain-induced martensitic transition [transformation-induced plasticity (TRIP)], generating a high-density dislocation network that significantly strengthens the material’s yield strength. These unique microstructural characteristics not only enhance the yield strength of the alloy but also improve its ductility through strain-induced phase transformation strengthening.

These findings illustrate that remarkable microstructural control is achievable in hierarchically heterogeneous alloys through advanced laser-based techniques. The intricate interplay between grain refinement, phase transformations, nanoprecipitates, and the formation of crystalline-amorphous heterostructures are expected to contribute to better mechanical performance, paving the way for the development of advanced structural materials with superior properties.

Strength-ductility synergy

For conventional metals and alloys, strength and ductility are mutually exclusive, as illustrated by the curve in Figure 12. Overcoming this strength-ductility trade-off is a major challenge in metallic materials^[57,128]. Reducing the grain size to nanoscale is an effective strategy to increase the strength, as grain boundaries can effectively block dislocation motion. However, this grain refinement also leads to a significant loss in ductility due to the constraint of plastic deformation mechanisms within small grains. Heterostructured materials, however, provide a novel pathway to address this dilemma^[129,130]. Extensive experimental studies have demonstrated that these structured materials possess an exceptional combination of mechanical properties, often overcoming traditional trade-offs. As illustrated in Figure 12, various microstructural designs created by laser-based techniques demonstrate exceptional mechanical properties that often exceed the traditional limitations seen in homogeneous MPEAs. The unique configurations of heterostructured materials represented by different marks illustrate that the laser-processed variants can achieve higher yield strengths without a proportional loss in elongation, in contrast to conventional homogeneous MPEAs (depicted as purple squares) that tend to cluster in the lower strength and ductility regions. Independent heterostructures, such as lamellar configurations or nanoprecipitates, demonstrate significant benefits for achieving strength enhancements but less improvement in ductility. However, core-shell structures and hierarchically heterogeneous structures provide notable advantages in plasticity retention. Consequently, the combination of multiple hetrostructures may inspire us more to solve the challenge of strength-ductility trade-offs.

Figure 12. Normalized yield strength versus uniform elongation of homogeneous and laser-processed heterostructured MPEAs^{[28,35,86,88,90,95,97-100,102,103-127]}.

Here are some typical examples of how specific heterostructures affect their performance. Recent investigations by Zhang et al. have revealed how inverse gradient-grained CoCrFeMnNi high-entropy alloys effectively transcend the conventional trade-off between strength and ductility^[88]. Through a novel processing method involving cold rolling followed by laser surface heat treatment, these alloys exhibit a unique gradient in grain sizes, which transitions from smaller and harder grains at the core to larger and softer grains at the surface. This gradation aligns closely with their microhardness profiles, optimizing the material’s mechanical properties.

This structural gradient not only enhances strength via HDI strengthening and strain hardening but also preserves ductility through the formation of high-order hierarchical nanotwins that interact synergistically with dislocations. The heterostructured samples outperform conventional homogeneous samples in tensile tests, exhibiting high yield strength (approximately 678 MPa), tensile strength (approximately 830 MPa), and notable ductility (about 28.2%).

The inverse gradient grain structure plays a pivotal role during tensile testing by promoting the formation of geometrically necessary dislocations and deformation twins. These features accommodate the strain gradient, significantly improving mechanical performance. Additionally, the interfaces between coarse and fine grains serve as sites for accumulating high densities of dislocations, generating back stresses that restrict dislocation movement in the softer regions, thus contributing to enhanced yield strength and further work hardening.

Building on these foundational insights, Kim et al. have further expanded the application of inverse-gradient structures in CoCrFeMnNi MPEAs through a similar approach^[123]. Their method involved cold rolling followed by LSR on both sides of the alloy sheets, resulting in a multi-scale heterogeneous microstructure. This structure features a gradient, transitioning from coarse grains at the surface to fine grains in the middle, with partially recrystallized grains at the core, resembling Figure 4E. The inverse gradient heterostructured samples exhibit superior tensile properties, with a yield strength of approximately 907.6 MPa and an ultimate tensile strength of around 949.4 MPa. Unlike the previously mentioned outer-soft and inner-hard structures, these inverse gradient heterostructured samples demonstrate excellent bendability, as the outer coarse-grained region prevents crack initiation. Although the presence of a coarse-grained area may reduce tensile performance compared to other gradient materials, it still outperforms homogeneous materials, and the decrease in strength differences between neighboring domains mitigates damage evolution. Thus, the produced heterogeneous HEAs exhibit superior tensile and forming properties.

Building upon the previously discussed strength-ductility trade-off, heterostructured materials have emerged as a promising solution to enhance work hardening capacity as well^[131,132]. This section delves into the work hardening phenomena in heterostructured materials, revealing their unique micro-mechanisms and macroscopic manifestations.

Fu et al. investigated the strain distribution and work hardening behavior of CoCrFeNiMn_0.75Cu_0.25 MPEA treated by LSP under uniaxial tensile loading, and the gradient structures formed inside the component are similar to those illustrated in Figure 3^[122]. It was found that the yield strength of the sample after four LSPs was increased by 1.5 times, and the elongation reached 40%. LSP treatment significantly improves the strain hardening ability of MPEA, which is mainly achieved by the combination of dislocation hardening and mechanical twins. The refined nanograins and nanotwins in the gradient structure form a dense shear band at the low strain stage and remain stable throughout the plastic deformation process, thus accommodating the major tensile strain and improving the overall strain hardening effect. In addition, the multi-axial strain transformation in the graded structural materials promotes the activation of multiple slip systems, further enhancing dislocation hardening. Consequently, this results in the enhancement of both strength and hardness while maintaining good plasticity in the interior, achieving a synergistic growth of strength and ductility. It should be noted that, for materials with internal porosity or cracks, cracks typically initiate at the weakest points surrounded by tensile stress during tensile tests, subsequently propagating until fracture occurs. However, shock waves influence the target during the LSP process, compressing porosity and transforming the stress state around these regions from tensile to compressive, significantly improving tensile performance.

Recent advancements by Zhu et al. have demonstrated the significant work hardening capacity in an interstitial solute-strengthened MPEA, Fe_49.5Mn₃₀Co₁₀Cr₁₀C_0.5, fabricated via SLM^[115]. The as-built MPEA exhibits a hierarchically heterogeneous microstructure, which is critical for its enhanced mechanical properties. The alloy boasts a high yield strength of approximately 710 MPa and an ultimate tensile strength of around 1 GPa, combined with a notable uniform elongation of about 28%, showcasing exceptional strength-ductility synergy.

The substantial work hardening in this MPEA is attributed to deformation mechanisms such as dislocation slip, deformation twinning, and phase transformation. Initially, high dislocation density within cellular structures leads to significant dislocation storage. As deformation progresses, planar slip bands and stacking faults enhance work hardening. The formation of deformation twins and deformation-induced hexagonal close-packed phases further accommodate deformation and act as strong barriers to dislocation movement, maintaining a high work hardening rate. The hierarchically heterogeneous microstructure, spanning several length scales, plays a pivotal role in mechanical performance. This microstructural design enhances back stress and HDI hardening, promoting homogeneous deformation and resulting in superior mechanical properties.

Similarly, Fe-based MPEAs were fabricated using SLM by Park et al., where the SLM process introduces a microstructure with spatial and compositional heterogeneity^[102]. High dislocation density is the main contributing factor to the high strength. Solute segregation induced by SLM lowers the phase stability of the matrix and promotes strain-induced martensitic transformation, as shown in Figure 13A-G. The synergy of high dislocation density and strain-induced martensitic transformation enhances the alloy's strength and ductility.

Figure 13. (A) Local strain distribution maps on the tensile-fractured SLM-Fe₆₀Co₁₅Ni₁₅Cr₁₀ obtained by the DIC analysis; (B) Evolutions of local strain and BCC phase fraction measured by DIC and EBSD analysis; (C-G) EBSD phase maps showing the deformation-induced martensitic transformation as a function of plastic deformation^[102]. (H) The schematic illustration of heterostructured Al₃CoCrCuFeNi MPEA; The simulated BCC to FCC phase-transformation process with the increased strain: 2% (I), 5.5% (J), 8% (K), and 10% (L) at the strain rate of 1 × 10¹⁰ s^-1; The simulated BCC to HCP phase-transformation process with the increased strain: 5% (M), 7% (N), 8% (O), and 10% (P) at the strain rate of 1 × 10¹⁰ s^-1[133].

In addition to experimental studies, simulations have been conducted to uncover the underlying mechanisms responsible for the property enhancements in heterostructured MPEAs. For example, Peng et al. conducted molecular dynamic simulations to investigate the rate-dependent deformation behavior and multistage strain hardening mechanisms of GNS MPEAs^[133] [Figure 13H-P]. Their simulations show that below a strain rate of 10⁸ s^-1, dislocation slip and deformation twinning are the main mechanisms. At a strain rate of 10⁹ s^-1, deformation twinning coexists with phase transformations, while at a strain rate of 10¹⁰ s^-1, phase transformations dominate. Additionally, edge dislocations, rather than screw dislocations, enhance the materials' strength. A microstructure-based constitutive model using a single parameter set is developed from the atomic simulations, accurately predicting the mechanical properties of MPEAs across various strain rates. Moreover, some structures can also be effectively produced through laser-based techniques, as demonstrated in Figure 13H. However, further investigations are still needed for specific systems prepared using different laser techniques. Challenges remain, particularly in obtaining precise parameters and in the complexity of modeling varying microstructures associated with laser-fabricated MPEAs.

These studies highlight a paradigm shift in understanding work hardening mechanisms in hetero-structured materials. Structural heterogeneity across multiple scales - from atomic to microscopic - creates a complex landscape for strain accommodation. This approach optimizes dislocation-microstructure interactions through gradient structures, nanotwins, multi-phases, and nanoprecipitates, leading to superior mechanical performance.

Fatigue resistance

Fatigue resistance is another critical consideration in the design and application of structural materials, as it directly influences the operational lifetime and safety of components subjected to cyclic loading. Interestingly, the unique microstructural architectures of heterostructure materials have been shown to significantly enhance the fatigue life of metallic systems^[134].

The study by Kim et al. highlights the exceptional capabilities of SLM in crafting CoCrFeMnNi MPEAs with refined microstructural features that substantially boost their high-cycle fatigue resistance^[117]. Unlike conventional manufacturing processes, SLM enables the precise control of microstructural evolution during the fabrication, leading to the formation of heterostructures characterized by nanosized oxide dispersions within a heterostructured metallic matrix. These oxides, formed in-situ during the SLM process, are strategically located at substructural and grain boundaries, serving as crucial barriers to dislocation movement and crack propagation.

In this heterostructured alloy, the fatigue limit is found to be significantly elevated at 570 MPa, a stark contrast to the 280 MPa observed in traditionally manufactured MPEAs [Figure 14A]. This improvement is not merely a function of enhanced material strength but is deeply rooted in the synergy between the nano-dispersed oxides and the underlying dislocation networks formed during the rapid solidification inherent to SLM. Additionally, the presence of deformation twins, induced under cyclic loading, plays a vital role in enhancing the material's ductility and in blunting potential crack tips, further impeding the fatigue crack propagation.

Figure 14. Fatigue resistance of MPEAs prepared by laser-based techniques: (A) S-N curves of SLM-CoCrFeMnNi^[117]. (B) Evolution of the yield stress at an offset strain of 0.01% after cyclic loading at different temperatures and strain rates of SLM-CoCrFeMnNi^[135]. High-cycle fatigue of SLM-CoCrFeNi: (C) S-N curve with the specimens used for the post-mortem examination indicated by the square box; (D) A plot of the measured maximum strain vs. fatigue cycles for the specimen tested under the maximum stress of σ_max= 450 MPa with a fatigue life of N_f = 1.06 × 105^[136].

Furthermore, the heterogeneous grain structure and dislocation networks typical of SLM-built materials contribute additional mechanical integrity by distributing stress more evenly across the material, thereby reducing the stress concentration at any single point. The structural complexity made by the SLM technique exemplifies the potential of heterostructured designs in pushing the boundaries of what is achievable in terms of mechanical performance and durability of MPEAs under cyclic loading conditions.

As the temperature changes, the performance sometimes might be largely different. Jin et al. examined the fatigue behavior of CoCrFeMnNi MPEAs fabricated by SLM across a temperature range of 22-600 °C^[135]. The study found that the alloy exhibits transient cyclic hardening followed by cyclic softening at all temperatures tested [Figure 14B]. The high dislocation density, resulting from the manufacturing process, effectively absorbs most applied strain, leading to initial hardening.

However, as temperature increases, the reduction in dislocation density results in decreased cyclic strength and related properties, such as yield strength and friction stress. Notably, the transition from transgranular to intergranular fracture morphology at elevated temperatures significantly influences fatigue life, highlighting the critical role of microstructural stability under varying conditions.

Chen et al. explored the microstructural evolution of a CoCrFeNi MPEA manufactured using SLM under high-cycle fatigue conditions^[136]. In their tests, deformation twinning is observed at a maximum stress of 450 MPa, while no twinning occurs at lower stress levels [Figure 14C and D]. This twinning leads to cyclic softening, accompanied by an increase in hardness, demonstrating the complex interplay between microstructure and fatigue behavior. The evolution of stacking faults and a maze-like dislocation structure within the alloy is another critical mechanism influencing fatigue performance. These features contribute to the alloy’s ability to accommodate strain and enhance resistance to crack initiation.

Those enhanced fatigue resistance observed in these heterostructured MPEAs can be attributed to several key mechanisms: Special microstructural features such as the presence of nanosized oxide dispersions and a heterogeneous grain structure improve resistance to crack initiation and propagation; The intricate dislocation networks formed during processing allow for effective energy dissipation during cyclic loading; The generation of deformation twins under stress not only improves ductility but also serves as a dynamic strengthening mechanism and the stress distribution; The heterogeneous microstructure facilitates uniform stress distribution, reducing localized stress concentrations that can lead to premature failure.

Currently, relatively few studies have explored the fatigue properties of heterostructured MPEAs. Understanding the underlying mechanisms is essential for optimizing alloy design and ensuring the longevity and safety of structural components in demanding applications, especially in extreme environments.

Wear resistance

Wear and friction significantly influence the service life and reliability of industrial components. With the advances in mechanical properties, heterostructured metallic materials have also shown substantial improvements in wear and friction resistance. The unique microstructural features and tailored interfaces play a vital role in modulating the tribological performance of these advanced MPEAs as well.

A notable example of this approach is demonstrated in a heterostructured TiZrHfTaNb MPEA fabricated via LSR^[93]. Luo et al. created a ~100 µm-thick layer with grain sizes reduced from ~200 µm to ~8 nm at the surface, increasing microhardness from ~240 to ~650 HV - a 2.7-fold improvement^[93]. This refinement also leads to phase decomposition and a depth-dependent phase distribution, significantly enhancing wear resistance by reducing wear rates by an order of magnitude compared to the base alloy. The researchers attributed this remarkable enhancement in tribological performance to the synergistic effects of the decomposed multi-phases and the gradient nanostructures, which collectively mitigated various wear mechanisms, including abrasive, oxidative, and adhesive wear. For systems prone to phase decomposition or the formation of amorphous structures, the selected LSR technique can induce a gradient distribution of solidification and cooling rates along the depth direction. This means the LSR process effectively melts the outermost surface region (resulting in an amorphous state), which will also experience the highest cooling rates. Consequently, crystallization during the subsequent solidification process may be only partially completed, leading to the formation of a crystalline-amorphous nanostructure. Additionally, rapid solidification processes (lasting only a few tens of milliseconds) can significantly inhibit atomic diffusion for element homogenization and the growth of newly formed second phases, particularly for sluggish elements. The rapid gradient cooling rates during the LSR process result in the growth of gradient grains in the newly formed phases, ultimately leading to the formation of a GNS layer^[93,101].

Another common approach is introducing new components for better surface properties; for instance, the recent research by Guan et al. introduces an innovative W/FeCoCrNi-based MPEA gradient coating, developed through LC^[92]. This study highlights the development of a dual-phase structure with Laves and face-centered cubic phases, engineered via the two-way diffusion behavior of FeCoCrNi and W elements during solidification. The resulting microstructure, characterized by a strategic hardness gradient and phase distribution, optimizes both hardness and plasticity. The hard Laves phase provides excellent wear resistance, while the ductile face-centered cubic phase maintains the coating’s plastic deformation capabilities, crucial for operational toughness.

Moreover, the coating's robust wear resistance is notably enhanced by a protective oxide film formed during the wear process. At elevated temperatures, the synergistic interaction between the hard and ductile phases effectively reduces stress concentrations during friction, optimizing the balance between hardness and ductility, essential for maintaining structural integrity under extreme conditions. Additionally, the transitional eutectic layer between these phases acts as a stress buffer, further enhancing the coating’s durability. Table 1 summarizes the wear characteristics of some homogenous and laser-fabricated MPEAs reported in recent years. It can be seen that most of the MPEAs with heterostructured layers from LC and LSR exhibit much lower wear rates and smaller friction coefficients. While the wear resistance of MPEA components by SLM is rarely reported, it also deserves to be well studied.