Enhanced strength and ductility in a high-entropy alloy via ordered oxygen complexes

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Abstract

Oxygen, one of the most abundant elements on Earth, often forms an undesired interstitial impurity or ceramic phase (such as an oxide particle) in metallic materials. Even when it adds strength, oxygen doping renders metals brittle1,2,3. Here we show that oxygen can take the form of ordered oxygen complexes, a state in between oxide particles and frequently occurring random interstitials. Unlike traditional interstitial strengthening4,5, such ordered interstitial complexes lead to unprecedented enhancement in both strength and ductility in compositionally complex solid solutions, the so-called high-entropy alloys (HEAs)6,7,8,9,10. The tensile strength is enhanced (by 48.5 ± 1.8 per cent) and ductility is substantially improved (by 95.2 ± 8.1 per cent) when doping a model TiZrHfNb HEA with 2.0 atomic per cent oxygen, thus breaking the long-standing strength–ductility trade-off11. The oxygen complexes are ordered nanoscale regions within the HEA characterized by (O, Zr, Ti)-rich atomic complexes whose formation is promoted by the existence of chemical short-range ordering among some of the substitutional matrix elements in the HEAs. Carbon has been reported to improve strength and ductility simultaneously in face-centred cubic HEAs12, by lowering the stacking fault energy and increasing the lattice friction stress. By contrast, the ordered interstitial complexes described here change the dislocation shear mode from planar slip to wavy slip, and promote double cross-slip and thus dislocation multiplication through the formation of Frank–Read sources (a mechanism explaining the generation of multiple dislocations) during deformation. This ordered interstitial complex-mediated strain-hardening mechanism should be particularly useful in Ti-, Zr- and Hf-containing alloys, in which interstitial elements are highly undesirable owing to their embrittlement effects, and in alloys where tuning the stacking fault energy and exploiting athermal transformations13 do not lead to property enhancement. These results provide insight into the role of interstitial solid solutions and associated ordering strengthening mechanisms in metallic materials.

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Fig. 1: Mechanical properties.
Fig. 2: Microscopic structure.
Fig. 3: Deformation mode.
Fig. 4: Intrinsic mechanism.

Data availability

The data that support the findings of this study are available from the corresponding authors on reasonable request.

Change history

  • 19 December 2018

    Change history: In this Letter, owing to a production error, all the data points (except the two points for O-2 and N-2, respectively) were missing in Fig. 1b. The figure has been corrected online.

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Acknowledgements

This research was supported by National Natural Science Foundation of China (grant numbers 51671018, 11790293, 51871016, 51531001 and 51671021), the 111 Project (grant number B07003), the Program for Changjiang Scholars and Innovative Research Team in University of China (grant number IRT_14R05) and the Projects of SKLAMM-USTB (grant numbers 2018Z-01 and 2018Z-19). Yuan W. acknowledges financial support from the Top-Notch Young Talents Program. Yuan W. and Hui W. acknowledges financial support from the Fundamental Research Funds for the Central Universities. We thank F. Zhang at the University of Science and Technology Beijing for help with synchrotron XRD. We also thank H. L. Huang at the University of Science and Technology Beijing and L. Qi and X. J. Zhao at the Chongqing University for help with TEM/STEM characterization and discussion.

Author information

Z. Lu designed the study. Z. Lei, X.L., Yuan W., Hui W., S.J., S.W., X.H. and Yidong W. carried out the main experiments. Z. Lei, X.L., Yuan W., Z. Lu, B.G. and D.R. analysed the data and wrote the main draft of the paper. L.G., Q.Z., H.C., Hongtao W. and J.L. conducted the TEM and STEM characterizations. B.G., P.K. and D.R. prepared the atom probe tomography specimens, processed the data and interpreted the results. K.A. conducted the neutron diffraction. Q.Z. conducted the synchrotron XRD. All authors contributed to the discussion of the results, and commented on the manuscript.

Correspondence to Zhaoping Lu.

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Extended data figures and tables

Extended Data Fig. 1 Aberration-corrected STEM of the as-cast HEAs.

Shown are the HAADF-STEM images for the [011]b.c.c. crystal axis with differently adjusted contrasts to show the existence of chemical short-range ordering, and the corresponding STEM-ABF images, for the equiatomic TiZrHfNb high-entropy base alloy (ac) and for N-2 HEA (that is, (TiZrHfNb)98N2) (df). The red panel represents the Zr/Ti-rich region, while the yellow panel indicates the Hf/Nb-rich region. No ordered interstitial occupation is observed in these two HEAs. Red squares represent the Zr/Ti-rich region and yellow squares indicate the Hf/Nb-rich region.

Extended Data Fig. 2 Occupation possibility analysis of interstitial oxygen/nitrogen from first-principles calculations.

a, Statistical distribution of system energy for the case of oxygen/nitrogen at different interstitial sites in the TiZrHfNb HEA. b, Comparison of average free energy for the systems with oxygen/nitrogen atoms at octahedral and tetrahedral interstitial sites. It can be seen that the octahedral interstitial oxygen/nitrogen has a free energy nearly identical to that of the tetrahedral interstitial oxygen/nitrogen, indicating that the likelihood of oxygen/nitrogen atoms occupying the tetrahedral or octahedral interstitial sites in the b.c.c. lattice is similar. The error bars represent the standard error of the mean.

Extended Data Fig. 3 Three-dimensional reconstruction of the O-2 HEA atom probe tomography dataset.

a, b, Randomized (a) and experimental (b) datasets on which an iso-composition surface encompassing regions in the point cloud containing more than 3.0 at% O was superimposed. The experimental dataset clearly shows evidence for OOCs.

Extended Data Fig. 4 Internal-friction measurements.

Internal-friction results obtained for the O-2 (that is, (TiZrHfNb)98O2) and N-2 (that is, (TiZrHfNb)98N2) HEAs. Metals containing solute atoms in interstitial solution show Snoek relaxation behaviour owing to stress-induced ordering43, which gives rise to a peak in the corresponding internal-friction spectrum (that is, the Snoek peak)63, and different internal-friction peaks correspond to different types of local atomic environments of the interstitials. Tan Delta represents the damping capacity. a, For the oxygen-doped alloy (TiZrHfNb)98O2 two peaks are observed: a dominant high-temperature peak and an additional low-temperature peak. b, For the nitrogen-doped alloy (TiZrHfNb)98N2 only the main peak is observed. This observation suggests that addition of oxygen to the TiZrHfNb HEA induces formation of two different types of interstitial atomic structures, unlike in the N-2 HEA, where only a single solid-solution peak appears.

Extended Data Fig. 5 X-ray and in situ neutron diffraction measurements.

a, XRD patterns of the TiZrHfNb base alloy, O-2 HEA (that is, (TiZrHfNb)98O2) and N-2 HEA (that is, (TiZrHfNb)98N2) with different pre-tensioned strains. bd, In situ neutron diffraction patterns of the three alloys. d is the interplanar distance. Both ex situ XRD and in situ neutron diffraction measurements confirm that there is no phase transformation in the three HEAs during deformation.

Extended Data Fig. 6 Transmission electron microscopy.

ac, TEM images of the as-cast equiatomic TiZrHfNb base alloy, O-2 HEA (that is, (TiZrHfNb)98O2) and N-2 HEA (that is, (TiZrHfNb)98N2). df, TEM images of the fractured HEA specimens. The TEM results further confirm that no second phase appears before and after the tensile tests. The inset in each figure is the corresponding electron diffraction pattern of the selected area.

Extended Data Fig. 7 Dislocation configuration.

a, For the equiatomic TiZrHfNb base alloy, at low tensile strain (2.5% strain), dislocations in linear arrays are observed. As the strain increases to 8%, planar slip bands and individual dislocation-rich sheets are formed. After fracture, although irregular dislocation cells can be seen, there exist several microbands, indicating that planar slip is still the dominant deformation mode. b, For the oxygen-doped alloy variant O-2 HEA (that is, (TiZrHfNb)98O2) at 2.5% strain, however, the dislocations are arranged in bundles and loops. At 8% strain, dislocation walls are formed. For the dislocation substructure after fracture, dipolar walls that mainly contain primary dislocation dipoles at a high density are observed, suggesting a typical cell-forming deformation microstructure in the O-2 HEA. c, For the nitrogen-doped alloy N-2 HEA (that is, (TiZrHfNb)98N2), the deformation mode is similar to that of the base alloy. In addition, slip traces at each specimen surface during deformation are also observed. Even at a low strain (2.5% strain), wavy slip lines are clearly seen in the oxygen doped variant O-2 HEA, whereas in the TiZrHfNb base alloy, even at high strain (8%), straight slip lines prevail and wavy slip lines only occur upon necking. Moreover, premature and much more serious necking occurs in the TiZrHfNb base alloy. It is worth mentioning here that intergranular fracture is observed in the nitrogen-doped variant N-2 HEA, which is probably caused by grain boundary segregation of nitrogen.

Extended Data Fig. 8 Pinning effect.

Aberration-corrected STEM observation of O-2 HEA (that is, (TiZrHfNb)98O2) after being pre-strained to 8%. B is the beam direction. The red arrows indicate the distinct dislocation pinning effect, which suppresses dislocation motion substantially during deformation.

Extended Data Fig. 9 Schematic diagram illustrating the plastic deformation mechanism in the oxygen-rich alloy

variant O-2 HEA.

Extended Data Fig. 10 Lattice parameter calculation.

ac, Plot of measured values of the lattice parameter versus \(\frac{{cos}^{2}{\vartheta }}{2}\left(\frac{1}{sin{\vartheta }}+\frac{1}{{\vartheta }}\right)\) for TiZrHfNb, (TiZrHfNb)98O2 and (TiZrHfNb)98N2 alloys. The position of each peak was measured on the diffractogram from which the lattice parameter was calculated. The measured lattice parameters were plotted versus \(\frac{{cos}^{2}{\vartheta }}{2}\left(\frac{1}{sin{\vartheta }}+\frac{1}{{\vartheta }}\right)\), where ϑ is the Bragg angle for each peak and the resulting graph extrapolated to zero to obtain the best value of the lattice parameter64. d, The calculated lattice parameters of the three alloys.

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