2023 Volume 64 Issue 11 Pages 2591-2595
One challenge hampering the structural applications of High-entropy alloys (HEAs) is to overcome the strength-ductility trade-off. Inspired by nature, engineering the HEAs with laminated structure composed of alternating soft and hard layers has the potential to strengthen the HEAs while maintaining the ductility. Here, taking the AlxCrMnFeCoNi HEA as the model material, the laminated HEAs composed by soft and hard layers were successfully prepared by laser additive manufacturing (LAM). The soft and hard laminated layers were achieved through tailoring the Al content. In several combinations of hard and soft, the strength of Al10–Al12 was significantly higher than that of Al12 HEA, indicating that an appropriate soft and hard matching can indeed further strengthen the HEA compared with the homogeneous HEA.
High-entropy alloys (HEAs), a new type of equiatomic or near equiatomic multicomponent alloys, have received considerable attention due to their excellent structural stability and exceptional resistances to fatigue, creep and corrosion.1–3) One challenge hampering the structural applications of HEAs, however, is to overcome the strength-ductility trade-off, which is also the perennially problematic in the materials science community.4–6) Specifically, the face-centered-cubic (fcc) HEAs (e.g. CrMnFeCoNi alloy) are ductile at room temperature, but typically not strong.7–9) In contrast, the body-centered-cubic (bcc) HEAs (e.g. AlCoCrFeNiTi0.5) are strong, but usually exhibit poor plasticity.10,11) A few traditional technical routes have demonstrated the capability to evade this dilemma in HEAs, such as by introducing secondary phases and controlling their sizes, morphologies and distribution through metallurgical processes and subsequent heat treatments,12–15) or achieving heterogeneous structure via cold rolling and subsequent annealing.16,17) Nevertheless, to improve the strength while retain the ductility of final parts with complex shapes from these methods requires many subsequent additional machining. It is therefore urgent to adopt a new technology that can realize the shaping of complex parts, while improve their strength simultaneously.
Over the eons of evolution, natural materials have developed a variety of delicate structure with exceptional properties and can provide excellent inspiration for the design and fabrication of novel materials.18,19) In the quest for strong and ductile solutions for structural materials, nacre, with the laminated structure and unique combination of remarkable strength and toughness, has received great interest as the model system.20–23) For instance, through mimicking the nacreous layer, assembling soft coarse grain layers and hard ultrafine grain layers alternatively can produce stronger and tougher metals.24–26) The hard layers serve as the primary load bearer for strength. The soft layers maintain good ductility by suppressing strain localization and dissipating energy in the hard layers through effective interface confinement.27–29) Therefore, engineering the HEAs with laminated structure composed of alternating soft and hard layers has the potential to strengthen the mechanical properties of HEAs.
Traditional approaches for the fabrication of laminated metals, like pressing and sintering or repeated rolling, also require intense tooling or mechanical treatment and are not readily applicable to the complex geometry components necessary for practical applications. Recent advances in laser additive manufacturing (LAM) technology make the fabrication of metallic components with arbitrary geometry possible and have promoted a surge in the investigation of biomimetic laminated metals.3,30–33) Unlike conventional manufacturing/processing techniques, LAM ‘prints’ final metallic parts directly from computer-aided design files, and thus offers the unique advantages of design freedom for complex geometry without the need for subsequent tooling. The LAM process is also characterized by highly localized melting and solidification processes, and hence generates unique microstructures that cannot be easily accessible via traditional metallurgical routes. More importantly, since the LAM is a powder-based layer-by-layer shaping and consolidation process, it permits ample opportunities to in-situ tailor the composition and microstructure of each layer during the deposition process, making it especially ideal for fabricating laminated HEAs.
In the present study, employing the AlxCrMnFeCoNi HEA as the model material, we fabricated laminated HEAs constructed by soft and hard layers to enhance the yield strength while maintaining the original tensile ductility. The soft and hard laminated layers were achieved through tailoring the Al content of the HEAs by accurately and instantaneously adjusting the powder feeding speed during the LAM process. The results showed that the appropriate soft and hard matching can indeed further strengthen the HEA compared with the homogeneous HEA. Our experiment provides a feasible method to improve the comprehensive mechanical properties of HEAs, and the introduced LAM technique enables the potential structural applications of HEAs at large scale.
The spherical pre-alloyed CrMnFeCoNi and elemental Al powders prepared by nitrogen gas atomization were employed as the raw materials. The CrMnFeCoNi HEA and Al powders are in the same size range of 40∼120 µm. The scanning electron microscope (SEM) morphologies of the two powders are shown in Fig. 1(a). LAM experiments were performed using the coaxial powder feed laser cladding system supplied with a 6 kW fiber laser, as illustrated in Fig. 1(b). To prepare the laminated HEAs, the CrMnFeCoNi and Al powders, stored separately in two powder hoppers, were fed through four coaxial nozzles by argon flow and delivered simultaneously into the melt pool created by the laser beam. Figure 1(c) shows that the laser printing method in this experiment is double-layer superposition. The purpose of double-layer is to avoid the uneven composition of the previous layer during the cladding of the next layer during laser processing, and to reserve sufficient processing allowance for the subsequent cutting of the tensile sample. If the soft and hard layer is designed too thick, the strengthening effect of the interface is not so obvious. The laser beam was focused on the substrate surface to a spot of ∼2.5 mm in diameter. The CrMnFeCoNi HEA and Al powders were injected from a hopper into the melt pool by argon flow through four coaxial nozzles. LAM experiments were conducted inside a working chamber filled with argon gas to keep the oxygen level less than 10 ppm. Meanwhile, argon was used as the shielding gas for the laser head. To obtain a continuous deposition, the overlapping fraction of the parallel tracks is set as 30%. During multi-layer deposition process, the layer thickness in the build direction was set as 0.5 mm. 45# steel plate was used as the substrate due to its high yield strength and comparable linear expansion coefficient with the AlxCrMnFeCoNi HEA.
(a) SEM of the gas-atomized CrMnFeCoNi and Al powders, (b) A schematic diagram of the LAM experiment, (c) Schematic diagram of high entropy alloy with double layer structure, (d) The outer appearance of a typical LAMed AlxCrMnFeCoNi HEA.
The microstructure and the chemical composition of the LAMed laminated AlxCrMnFeCoNi HEA were studied by a Zeiss Ultra 55 field-emission scanning electron microscope (SEM) equipped with electron backscatter diffraction (EBSD). The phases of the LAMed AlxCrMnFeCoNi HEA were characterized by X-ray diffraction (XRD, Empyrean).
2.3 Mechanical testingThe dog-bone-shaped tensile specimens with the gauge dimension of 16 mm (length) × 3 mm (width) × 3 mm (thickness) were cut from the LAMed samples by electrical discharge machining. Tensile tests were conducted on a SANS CMT5305 testing machine at room temperature under a strain rate of 1 × 10−4 s−1. The tensile strain was calculated as the total elongation (measured by a standard extensometer) with respect to the initial gauge length. Each sample condition was repeated three times. In order to observe the microstructural evolution upon deformation, the surfaces of the tensile specimens were carefully polished to a mirror finish prior to tension.
The laminated structure for the laminated HEAs is designed as overlaying layers of hard and soft, as schematically shown in Fig. 1(c). The thickness of each layer is 0.5 mm. In order to improve the strength without losing the ductility of HEA, we tailor the mechanical properties of the soft and hard layers to achieve the best combination of soft and hard. This can be readily achieved by alloying Al element into the CrMnFeCoNi HEA. Many previous studies have found that the slight addition of Al can significantly change the mechanical properties of the CrMnFeCoNi HEA through altering the lattice stability and precipitation of the hard phase.14) This high sensitivity of the mechanical property of the CrMnFeCoNi HEA to the Al element is leveraged in this study to design the laminated structure. Along this technical route, we fixed the Al content of one layer at 12 at%, and then adjusted the Al content of the other layer, as illustrated in Fig. 1(c).
To accurately quantify the relationship between the Al content and mechanical property, we firstly synthesized a serious of rectangular AlxCrMnFeCoNi HEAs (denoted as Alx HEA hereafter) with Al content ranging over x = 0∼12 at% by LAM. A typical outer appearance of the LAMed HEA is shown in Fig. 1(d). A summary of the XRD patterns for the AlxCrMnFeCoNi HEAs is shown in Fig. 3(a). Apparently, the AlxCrMnFeCoNi HEAs undergo a phase structure transition from fcc to bcc as the Al content increases from 0 to 12 at%. When the Al content is no more than 6 at%, the XRD patterns reveal a single fcc phase, indicating that the addition of a small amount of Al does not cause the conversion from fcc to bcc. When the Al content increases to 8 at%, a small (110)bcc peak starts to appear near the predominant (111)fcc peak, demonstrating that the original phase constitution changes from single fcc phase to fcc + bcc dual phases. Afterwards, the intensity of the (110)bcc diffraction peak becomes stronger as the Al content further increases to 12 at%, indicating the increment in the volume fraction of bcc phase. A typical TEM bright field image and corresponding selected-area electron diffraction pattern of Al10 HEA is presented in Fig. 2 to verify the bcc and fcc composite structure.
TEM bright field image and the selected-area electron diffraction pattern of Al10 HEA.
To further verify the effect of Al element on the microstructure, we collected the EBSD data of the AlxCrMnFeCoNi HEAs. Figures 3(b)–(e) exhibit the EBSD phase maps for the Al0, Al8, Al10 and Al12 HEAs. Clearly, the phase maps of Al0 and Al8 HEAs presented in Figs. 3(c) and (d) confirm that the fcc phase (represented by blue color) is the single and near-single phase in these samples. When the Al content gradually increases to 8 at%, a small amount of bcc phase precipitation (represented by red color) is distributed on the blue fcc matrix, as shown in Fig. 3(d). As the Al content increases to 10 at%, a large amount of bcc phase is dispersed, as shown in Fig. 3(e). By using the software ImageJ to estimate the volume fraction of bcc phase through the EBSD images, the contents of the bcc phase are 0%, 7%, 26% and 71% for the Al0, Al8, Al10 and Al12 HEAs, respectively. In addition, Figs. 3(c) and (d) show bcc phase precipitating from the fcc matrix, this feature will affect the growth of columnar grain direction and size. Obviously, when the Al content is increased to a certain extent, there will be more equiaxed grains in the alloy. This is because the increase of Al content will lead to more bcc phase precipitation, and the bcc phase growing will affect the growth process of columnar crystals to a certain extent, thus changing the grain morphology.
(a) XRD patterns of the AlxCrMnFeCoNi HEAs with different Al contents; (b), (c), (d), and (e) EBSD phase maps of the Al0, Al8, Al10 and Al12 HEAs. The fcc and bcc phases are respectively represented by blue and red colors in the phase map images.
The strength of AlxCrMnFeCoNi HEAs increases with the increase of bcc phase. Therefore, we tested the tensile properties of Al12 HEA as a reference, and further improved the strength by regulating the combination of soft and hard of laminated HEAs. The tested room temperature engineering stress-strain curve of the Al12 HEA is presented in Fig. 4.
Room temperature engineering stress-strain curves of the two-lamella HEAs.
The macroscopic mechanical properties of the laminated HEA do not follow the rule of mixture (ROM) by simply summing the mechanical properties of the soft and hard lamellae proportionally, because there is interaction between adjacent soft and hard lamellae with different mechanical properties. The interaction between adjacent soft and hard lamellae, also referred as “interface coupling effect”, provide supererogatory strengthening and extraordinary strain hardening behavior beyond the ROM during the interface-coordinated deformation.28) Then the mechanical properties of the laminated HEA are the sum of the mechanical properties of the soft lamellae, the hard lamellae and the interface coupling effect between them. Therefore, it is possible to maximize the contribution of the interface coupling effect through optimizing the mechanical mismatch between the soft and hard layers. In order to screen out the best match of hard and soft lamellae, we rapidly produced a series of two-lamella HEAs by LAM. One lamella of the two-lamella structure is the hardest HEA containing Al12, and the other one is a relatively soft HEA with less Al. For instance, the two-lamella structure composed by soft Al4 HEA and hard Al12 HEA is labeled as Al4–Al12 hereafter. The tensile stress-strain curves of these two-lamella HEAs are shown in Fig. 4. It is obvious that the yield strength of the two-lamella HEAs gradually increases from the Al4–Al12 to Al10–Al12, while the tensile elongation of the two-lamella structure changes little and fluctuates in the range of 5%–5.5%. It is noteworthy that compared with Al12 HEA, the three combinations Al4–Al12, Al6–Al12 and Al8–Al12 showed no advantage in strength, indicating that although the combination of the two layers that differ greatly in hardness and softness can achieve a greater “interface coupling effect”, the softer layer will also significantly reduce the overall strength.22) The strength of Al10–Al12 was significantly higher than that of Al12 HEA, proving that adding soft and hard matching structure into homogeneous HEA can indeed further strengthen the HEA. The yield strength and elongation of the best combination Al10–Al12 are 650 MPa and 5.5%, respectively.
The microstructure difference between the hard and soft layers can lead to discordant deformation near the interfaces, materials in the process of tensile deformation, the soft zone started plastic deformation, and hard area to maintain flexibility, soft back stress to offset some of the shear stress in the area.27,29) The presence of this back stress can enhance laminated HEA. To reveal the microstructures at the interfaces of the laminated HEAs, the EBSD inverse pole figure (IPF) and phase maps for the laminated HEAs were collected as shown in Fig. 5. The sample locations are perpendicular to the laser scanning direction. Clearly, the soft lamella is almost a completely single fcc phase when the Al content less than 6%. As the Al content increases to 10%, the bcc phase precipitated on fcc matrix gradually increases. While within the hard lamella, bcc phase almost covers the whole matrix. It is worth noting that there is a transition zone between the soft and the hard lamella. The previous layer is partially melted during the printing process of the latter layer, causing the upper layer of Al to enter the next layer, thus changing the phase composition. In addition, within the Al4, Al6, Al8, Al10 lamella, the fine equiaxed grains tend to form on the top of the molten pools and coarse columnar grains distribute at the rest region. In contrast, there is no obvious difference in grain morphology within the Al12 lamella, which are all equiaxed grains.
(a), (b), (c), and (d) EBSD inverse pole figure (IPF) and phase maps of the laminated Al4–Al12, Al6–Al12, Al8–Al12, Al10–Al12 HEAs, the fcc and bcc phases are respectively represented by blue and red colors in the phase map images.
In summary, taking the AlxCrMnFeCoNi HEA as the model material, the laminated HEAs composed by soft and hard layers were successfully prepared by LAM. The soft and hard laminated layers were achieved through tailoring the Al content of the HEAs by accurately and instantaneously adjusting the powder feeding speed during the LAM process. In several combinations of hard and soft, the strength of Al10–Al12 was significantly higher than that of Al12 HEA, indicating that the mechanical properties of the laminated HEA are the sum of the mechanical properties of the soft lamellae, the hard lamellae and the interface coupling effect between them. The appropriate soft and hard matching can indeed further strengthen the HEA compared with the homogeneous HEA.
This work was supported by the National Natural Science Foundation of China under Grant No. 52271023.
