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Strong yet ductile nanolamellar high-entropy alloys by additive manufacturing

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Additive manufacturing produces net-shaped components layer by layer for engineering applications1,2,3,4,5,6,7. The additive manufacture of metal alloys by laser powder bed fusion (L-PBF) involves large temperature gradients and rapid cooling2,6, which enables microstructural refinement at the nanoscale to achieve high strength. However, high-strength nanostructured alloys produced by laser additive manufacturing often have limited ductility3. Here we use L-PBF to print dual-phase nanolamellar high-entropy alloys (HEAs) of AlCoCrFeNi2.1 that exhibit a combination of a high yield strength of about 1.3 gigapascals and a large uniform elongation of about 14 per cent, which surpasses those of other state-of-the-art additively manufactured metal alloys. The high yield strength stems from the strong strengthening effects of the dual-phase structures that consist of alternating face-centred cubic and body-centred cubic nanolamellae; the body-centred cubic nanolamellae exhibit higher strengths and higher hardening rates than the face-centred cubic nanolamellae. The large tensile ductility arises owing to the high work-hardening capability of the as-printed hierarchical microstructures in the form of dual-phase nanolamellae embedded in microscale eutectic colonies, which have nearly random orientations to promote isotropic mechanical properties. The mechanistic insights into the deformation behaviour of additively manufactured HEAs have broad implications for the development of hierarchical, dual- and multi-phase, nanostructured alloys with exceptional mechanical properties.

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Fig. 1: Microstructure of AM AlCoCrFeNi2.1 EHEA.
Fig. 2: Tensile properties of AM AlCoCrFeNi2.1 EHEAs.
Fig. 3: Lattice strains and stress partitioning in fcc and bcc phases during uniaxial tension.
Fig. 4: Meso- and atomic-scale deformation structures.

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Data availability

The data of this study are included in the article, the Extended Data and the Supplementary Information.

Code availability

The code used for finite-element analyses is publicly available on the GitHub repository at


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We thank D. Follette, P. Hou, M. Wu, K. A. Beyer and M. J. Frost for their experimental assistance. W.C. acknowledges support from the US National Science Foundation (DMR-2004429) and UMass Amherst Faculty Startup Fund. T.Z. acknowledges support from the US National Science Foundation (DMR-1810720 and DMR-2004412). Y.M.W. acknowledges support from the US National Science Foundation (DMR-2104933). T.V. acknowledges support from the Laboratory Directed Research and Development (LDRD) programme (21-LW-027) at Lawrence Livermore National Laboratory (LLNL). His work was performed under the auspices of the US Department of Energy (DOE) by LLNL under contract no. DE-AC52-07NA27344. In situ neutron-diffraction work was carried out at the Spallation Neutron Source (SNS), which is a US DOE user facility at the Oak Ridge National Laboratory (ORNL), sponsored by the Scientific User Facilities Division, Office of Basic Energy Sciences. APT research was supported by the Center for Nanophase Materials Sciences (CNMS), which is a US DOE Office of Science User Facility at ORNL. We thank J. Burns for assistance in performing the APT sample preparation and running the APT experiments. This research also used high-energy X-ray resources of the Advanced Photon Source (Beamline 11-ID-C), a US DOE Office of Science User Facility operated at Argonne National Laboratory under contract number DE-AC02-06CH11357.

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Authors and Affiliations



J.R. and W.C. developed the three-dimensional printing-process map. J.R. and F.K. fabricated all samples and performed the processing parameters optimization. J.R., Y.L., L.L. and S.P. performed the optical microscopy and SEM microstructure characterization and mechanical testing. D.Z., K.Y.X., G.G., T.V. and Y.M.W. performed the EBSD and TEM characterization and analyses. J.R., Y.C., K.A. and W.C. conducted in situ neutron-diffraction experiments and analysed the data. J.D.P. collected and analysed the APT data. S.G. conducted the thermodynamic calculation. Y.Z. and T.Z. developed the DP-CPFE model and performed numerical simulations. J.R., Y.Z., D.Z., K.Y.X. T.Z. and W.C. drafted the initial manuscript. W.C. conceived, designed and led the project. All co-authors contributed to the data analysis and discussion.

Corresponding authors

Correspondence to Ting Zhu or Wen Chen.

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Nature thanks Sheng Guo, Minh-Son Pham and the other, anonymous, reviewer(s) for their contribution to the peer review of this work.

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Extended data figures and tables

Extended Data Fig. 1 Pole figures of as-printed AlCoCrFeNi2.1 acquired by neutron diffraction.

a, Pole figures of FCC- (111), (200), (220), and (311) before loading. b, Pole figures of BCC-(110), (200), (211), and (321) before loading. c, Pole figures of FCC- (111), (200), (220), and (311) after fracture. Because the BCC peaks display extensive broadening after fracture, single-peak fittings are not convergent at lots of beam incident directions and pole figures of BCC orientations after fracture are not available. In all pole figures, the loading direction (LD) is out of plane, the transverse direction (TD) is along the horizontal direction, and the build direction (BD) is along the vertical direction. Before loading, the as-printed sample shows a rather weak texture with slightly preferred orientation of FCC-(110)//BD. After fracture, the FCC-(111)//LD texture is developed, suggesting prominent dislocation slips on {111} planes in the FCC phase.

Extended Data Fig. 2 High-resolution TEM image showing the consistent crystal structure within BCC nanolamellae.

The inset shows the corresponding Fast Fourier Transform (FFT) diffractogram of the entire area that can provide chemical ordering information. No alternating intensity variation is observed in the FFT diffractogram, suggesting that no apparent ordered B2 phase is present.

Extended Data Fig. 3 Extreme processing conditions enabled by L-PBF and the resulting highly metastable microstructure of multi-component eutectic alloys.

a, Comparison of cooling rate and thermal gradient between several additive manufacturing methods such as laser powder bed fusion (L-PBF) – used in this work, laser directed energy deposition (L-DED), wire arc additive manufacturing (WAAM), as well as conventional casting (CC) and directional solidification (DS)1. Extremely large cooling rates and thermal gradients are inherent to the unique spatial-temporal feature of L-PBF and thus give rise to the diffusion-limited solidification and far-from-equilibrium microstructure of our EHEAs. b, Schematic illustration of the cooling rate effects on microstructural morphologies and length scales for typical dual-phase multi-component eutectic alloys.

Extended Data Fig. 4 Kocks-Mecking plot showing the strain-hardening rate of as-printed AlCoCrFeNi2.1.

Strain-hardening rate (i.e., rate of increase of true stress with respect to true strain) is plotted as a function of true stress. Symbols represent experimental data points and the solid line is the fitting curve.

Extended Data Fig. 5 Tensile stress–strain curves of as-printed AlCoCrFeNi2.1 EHEAs along different directions.

Comparable mechanical properties of these samples at a similar build height demonstrate the isotropic mechanical behaviour of AM AlCoCrFeNi2.1 EHEA consisting of nanolamellar eutectic colonies with nearly random orientations.

Extended Data Fig. 6 TEM images showing stacking faults (SFs) in strained FCC nanolamellae.

a, SFs observed at the strain level of 5%. b, Same as a except at 15%. SFs are highlighted by yellow arrows.

Extended Data Fig. 7 Evolution of back stress during tensile deformation of as-printed AlCoCrFeNi2.1.

a, Loading–unloading-reloading (LUR) true stress–strain curve. b, A representative LUR cycle showing the hysteresis loop. The back stress is calculated by Dickson’s method and thus defined as σb = σ0σe = (σ0 + σu)/2 – σ*/2, where σb denotes the back stress, σ0 the flow stress before unloading, σe the effective stress, σu the unloading yield stress, and σ* the viscous stress. c, Flow stress, back stress, and effective stress versus true strain during tensile deformation. Error bars represent the standard deviation.

Extended Data Fig. 8 AM Ni40Co20Fe10Cr10Al18W2 EHEA with high strength and large tensile ductility.

a, 3D-reconstructed optical micrographs. b, Secondary electron micrograph showing the microscale eutectic colonies with different growth directions. c, Secondary electron micrograph revealing the typical nanolamellar structure. d, 3D-reconstructed EBSD IPF maps. The eutectic colony size distribution is obtained from the top-view map. The 001, 110, 111 pole figures of FCC phase are collected from the top-view EBSD map. Note that the BCC nanolamellae are difficult to index by EBSD due to their ultra-small thicknesses of ~35 nm. e, Lamellar thickness distribution of BCC and FCC lamellae in as-printed Ni40Co20Fe10Cr10Al18W2 EHEA. The average interlamellar spacing (λ ≈ 133 nm) is ~5 times smaller than that in the as-cast Ni40Co20Fe10Cr10Al18W2 (λ ≈ 0.82 μm). f, Neutron-diffraction pattern of AM Ni40Co20Fe10Cr10Al18W2 composed of FCC and BCC/B2 phases. g, Quasi-static tensile stress–strain curves of the as-cast and AM Ni40Co20Fe10Cr10Al18W2 EHEAs. Our AM EHEA exhibits a high yield strength of ~1.5 GPa and ultimate tensile strength of ~1.7 GPa, which outperform the as-cast counterpart by twofold with no significant loss in ductility. Note that the tensile stress–strain curve of the as-cast sample (dashed line) is taken from the literature; the substantially low elastic modulus and large elastic strain limit are likely due to the inaccurate strain measurement of this literature result.

Extended Data Table 1 Compositions of FCC and BCC phases in as-printed AlCoCrFeNi2.1, compared with as-cast counterpart25

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Ren, J., Zhang, Y., Zhao, D. et al. Strong yet ductile nanolamellar high-entropy alloys by additive manufacturing. Nature 608, 62–68 (2022).

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