Nanocrystalline High-Entropy Alloys: A New Paradigm in High

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Nanocrystalline high entropy alloys: A new paradigm in high temperature strength and stability Yu Zou, Jeffrey Martin Wheeler, Huan Ma, Philipp Okle, and Ralph Spolenak Nano Lett., Just Accepted Manuscript • DOI: 10.1021/acs.nanolett.6b04716 • Publication Date (Web): 26 Jan 2017 Downloaded from http://pubs.acs.org on January 30, 2017

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Nanocrystalline high entropy alloys: A new paradigm in high temperature strength and stability

Yu Zou1, *, Jeffrey M. Wheeler1, Huan Ma1, Philipp Okle1 and Ralph Spolenak1,* 1

Laboratory for Nanometallurgy, Department of Materials, ETH Zurich, Vladimir-Prelog-Weg 5, CH8093 Zurich, Switzerland

Abstract

Metals with nanometer scale grains, or nanocrystalline metals, exhibit high strengths at ambient conditions, yet their strengths substantially decrease with increasing temperature, rendering them unsuitable for usage at high temperatures. Here, we show that a nanocrystalline high entropy alloy (HEA) retains an extraordinarily high yield strength over 5 GPa up to 600 °C – one order of magnitude higher than that of its coarse-grained form, and five times higher than that of its single-crystalline equivalent. As a result, such nanostructured HEAs reveal strengthening figures of merit – normalized strength by the shear modulus above 1/50 and strength-to-density ratios above 0.4 MJ/kg, which are substantially higher than any previously reported values for nanocrystalline metals in the same homologous temperature range, as well as low strain-rate sensitivity of ~0.005. Nanocrystalline HEAs with these properties represent a new class of nanomaterials for high-stress and high-temperature applications in aerospace, civilian infrastructure and energy sectors.

Keywords: high entropy alloys; nanocrystalline; high temperature; strength; grain boundary; stability *Correspondence should be addressed to [email protected] (Y. Z., current address: Department of Mechanical Engineering, Massachusetts Institute of Technology, 77 Massachusetts Avenue, Cambridge, Massachusetts 02139, USA) or [email protected] (R. S.).

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For the past three decades, nanocrystalline (nc) metals have been attractive as potentially useful structural materials, primarily owing to their high strengths1-4. Their prominent strengths are attributed to a large volume fraction of grain boundaries as barriers to obstruct dislocation motion. According to the Hall-Petch law5, such nc-metals, down to the grain sizes of ~40 nm, can attain strength over an order of magnitude higher than that of coarse-grained ones. Grain boundaries, however, become a vulnerability for metals at high temperatures; for example, single-crystalline (sx) nickel-based superalloys are more preferred than polycrystalline ones to be used for gas turbine blades, because of significant grain-boundary diffusion and sliding in the latter at working temperatures6. In nc-metals, grain-boundary sliding7, diffusion-mediated plasticity8, and grain-boundary migration9 are even more pronounced, such that grains coarsen even at ambient or modest temperatures, and, consequently, material strength diminish10, 11. In general, nc-metals have never been considered to be used at elevated temperatures because of their poor high-temperature strengths and stability. Hence, developing nc-metallic systems that retain their high strengths over a wide temperature range has been highly desirable for both scientific interests and demanding applications.

Over the last few years, attention has been focused on solving the problem by two routes: grain boundary design and alloying design – the first introduces low energy grain boundaries12, for example, nanoscale twin boundaries13; the second stabilizes grain boundaries by elemental segregation14, 15. Using these methods, microstructural stability of these nc-metals has been improved for high-temperature, longduration conditions. The extant literature, however, reveals a large gap in demonstrating the mechanical stability of nc-metals under mechanical loading at a relatively high temperature, wherein grain-boundary motion and diffusion may occur and material strength is decreased. Very recently, Darling et al.16 reported that an immiscible Cu-Ta nc-metal showed high strength and creep resistance from room temperature to 600 °C, due to nanoscale Ta segregation. These studies have motivated us to achieve even stronger and more stable nc-metals by virtue of solid-solution strengthening and grain-boundary segregation.


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As an evolving field of metallurgy, high entropy alloys (HEAs) generally consist of five or more equiatomic elements and form a single solid-solution phase17, 18. A variety of interesting and useful properties emerge in face-centered cubic (fcc) coarse-grained HEAs including high fracture resistance19, high strengths20, and enhanced thermal stability21. Two typical body-centered cubic (bcc) refractory HEAs exhibit a yield strength of ~500-700 MPa at 1200 °C, even surpassing conventional nickel-based superalloys22. Mechanical properties of both fcc and bcc HEAs have been studied at macroscopic and microscopic levels23, 24. In such a chemically enriched system, multiple principle elements tend to strengthen materials markedly by the solid-solution hardening mechanism, meanwhile randomly distributed elements in grain interiors may lower relative grain-boundary energy, minimizing thermodynamic driving force for grain-boundary diffusion and migration. In addition, segregation may occur at grain boundaries and work as barriers to prevent grain-boundary motion. Our previous work demonstrates that, at room temperature, nc-HEAs (Nb25Ta25Mo25W25) exhibit higher strength than sx-HEA equivalents by a factor of two, due to grainboundary strengthening24, 25. Now, the question arises: Will the nc-HEAs retain their superior strength even at elevated temperatures, when their large volume fraction of grain boundaries becomes a manifest weakness?

Sx-HEA pillars were produced by focused ion beam milling from a 〈011〉-orientated coarse grain within a bulk HEA specimen. Nc-HEA pillars, which were milled from HEA films, consist of strongly textured columnar grains with an average grain size (diameter) of ~70-100 nm. We conducted compression tests on the sx- and nc-HEA micro-pillars (Supplementary Information Fig. S1) using a diamond flat punch over a temperature range from room temperature to 600 °C (Fig. 1). The sx-HEA pillars that are deformed at 25-400 °C show single slips with wavy morphologies (Figs. 1a-1c), while those deformed at 600 °C reveal sharper slip traces (Fig. 1d and Supplementary Information Fig. S2). After deformation, the nc-HEA pillars exhibit parallel wavy slip lines intersecting with grain boundaries after they were deformed from 25°C to 600 °C. The nc-HEA pillars that are deformed at 600 °C show more homogeneous deformation with fewer cracks, suggesting enhanced ductility at a higher test temperature (Fig. 1h). We also find that 3 ACS Paragon Plus Environment

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the grain sizes remain consistent (~70-100 nm) after deformation (Supplementary Information Fig. S2), as compared with those without deformation27. With increasing temperature, the yield strength of the sx-HEA pillars decrease from ~2 GPa at room temperature to ~1 GPa at 600 °C, with a relatively higher magnitude of stress drops (Fig. 1i); while the strength of nc-HEA pillars only decrease from ~7 GPa at room temperature to ~5 GPa at 600 °C, with a relatively consistent magnitude of stress drops (Fig. 1j).

Fig. 2a summarizes temperature dependence of yield strengths of the sx- and nc-HEA pillars. At 600 °C, the nc-HEA pillars still exhibit a high yield strength of ~5 GPa, which is nearly fivefold greater than that of the sx-HEAs and one order of magnitude higher than that of the bulk coarse-grained ones. The percentage of strength reduction as a function of temperature indicates that the strength softening behavior of a material is due to temperature increase. Fig. 2b shows that the bulk-HEAs22, sx-HEAs, and nc-Cu-Ta alloys16 exhibit a reduction of strength by ~50% at 600 °C, while the strength of the nc-HEAs is only reduced by ~20%.

Our nc-HEAs are attractive not only due to their higher absolute strength than other nc-metals (Supplementary information Fig. S3), but also because, in a wide range of temperatures, they offer the highest normalized strength (i.e., critical resolved shear stress, τ, divided by shear modulus, µ) and specific strength (i.e., strength-to-density ratio). Fig. 2c illustrates an Ashby-inspired map of the normalized strengths as a function of homologous temperatures (testing temperature, Tt, over melting temperature, Tm) for our sx- and nc-HEA pillars, as compared to the literature data for bulk coarse-grained bcc metals and nc-metals. With respect to compressive strength in the same homologous temperature range, the nc-HEAs significantly outperforms bulk bcc metals26, other nanostructured metals 27-32, high-temperature stable ncInconel 71833, and recently reported nc-Cu-Ta alloys16. Over the temperature range of 0.1-0.3 of Tm, the strength of our nc-HEAs falls in the ultra-strength regime as ~µ/50-µ/30, approaching its theoretical value of µ/10. In addition, Fig. 2d highlights the high strength-to-density ratios for the nc-HEAs as structural materials: our nc-HEAs achieve specific strength or elastic energy density of ~0.4-0.5 MJ/Kg, which markedly exceeds all nc-metals tested in the same temperature range. Such nc-HEAs exhibit comparable 4 ACS Paragon Plus Environment

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specific strength levels to that of carbon lattices and carbon-ceramic lattices (~0.2-0.8 MJ/kg)34, approaching the lower bound of diamond (~2 MJ/Kg). Although micro-pillar samples exhibit sizedependent strength26, such size effect in 1-µm refractory HEA pillars (twice strength of that of the bulk samples) are not as pronounced as that in fcc metal pillars 35.

To understand the observed thermal and mechanical stability of nc-HEAs, we pay attention to the microstructure and stress-strain curves again (Fig. 1). Stress drops during micro-compression are believed to be attributed to a dislocation avalanche phenomenon, in which dislocations nucleate, propagate, and escape free surfaces of sx-pillars in a short period of time. In nc-pillars, grain boundaries may prevent such dislocation avalanches occurring and decrease the magnitude of force drops. It is known that both the low magnitude of stress drops and wavy slip features indicate that screw dislocations have lower mobility than edge dislocations – a typical deformation behavior of bcc metals due to the cross slip of screw dislocations36. The high magnitude of stress drops and sharp slip bands of the sx-HEA pillars at 600 °C indicate that the mobility of screw dislocations is increased – a deformation feature similar to fcc metals. In addition, we utilized strain-rate jumps during the micro-compression test37 to determine the strain-rate sensitivity (m) of the flow stress of the sx- and nc-HEA pillars by measuring the flow stress, σ, as a function of strain rate, 𝜀̇: 𝑚=

𝜕 𝑙𝑛 𝜎 38 . 𝜕 𝑙𝑛 𝜀̇

The sx-HEA pillars exhibit decreasing strain-rate sensitivity with increasing temperature –

from ~0.025 at room temperature to ~0.01 at 600 °C, while the nc-HEA pillars show a consistent m value of ~0.005 (Fig. 3a).

Furthermore, according to the m values, we are able to calculate the apparent activation volume, Va, as 𝑉𝑎 = 𝐾 𝑇

𝐵 , 38 where kB is the Boltzmann constant of 8.617×10−5 eV/K and T is the absolute temperature. The √3 𝑚𝜎

value of Va is related to the area swept by dislocation segments during a single thermally activated event and, more importantly, is indicative of the deformation mechanism. Considering a Burgers vector of 3.229 Å24, we derive the Va for the sx- and nc-HEAs to be ~10 b3 at room temperature (Fig. 3b). This value is 5 ACS Paragon Plus Environment

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different from that of fcc metals, but identical to that reported from bulk tensile measurements on W, in which thermally activated double kink mechanism is believed to be responsible for the motion of screw dislocations39. With increasing temperature, the Va of the sc-HEA pillars increases to an average value of ~175 b3 at 600 °C, implying that sx-HEAs are deformed by, or partially by, the Peierls mechanism at 600 °C, which is comparable to the deformation of fcc metals28. The Va of the nc-HEAs slightly increases to ~50 b3 at 600 °C, suggesting the deformation of nc-HEAs could be still dominated by the kink-pair mechanism rather than the grain-boundary mediated mechanism. This is different from previous studies on fcc nanocrystalline nickel pillars, in which activation volumes increase with increasing test temperature, suggesting an enhanced grain-boundary mediated deformation37.

To further understand the role of the grain boundaries on the deformation of nc-HEAs at elevated temperatures, we used the atom probe tomography (APT) technique to characterize the local chemical arrangement of both sc- and nc-HEAs (also described in Ref [40]). They both show a uniform distribution of Nb, Mo, Ta and W without clustering (Figs. 4a and 4c), confirmed by one-dimensional (1D) concentration profiles (Figs. 4b and 4e). In the nc-HEA sample, we can clearly identify a planar region with enriched foreign elements (N, C, and O) from the top to bottom of the tip (see Supplementary information Fig. S4). This band feature with enriched N, C, and O can be correlated to a grain boundary, which is included in the nc-HEA tip and was observed using a transmission electron microscopy (TEM). At this grain boundary region, we also observe various nitrides and oxides (TaN, TaO NbN, NbO, and WN) with the average concentrations of between 1.6 at.-% and 0.05 at.-% (Figs. 4d and 4f). The foreign elements N, C, and O could be induced from raw materials, and the oxides and nitrides could be formed during sample preparation. Most importantly, although such nanoscale nitride and oxide clusters at grain boundaries may sacrifice material ductility, they are able to work as rigid interfaces to impede grain boundary motion, contributing a kinetic mechanism for enhancing microstructural stability at elevated temperatures. In addition, due to highly chemically disordered structures in both grain interiors and grain boundaries in the HEAs, the thermodynamic driving force for boundary motion is low, contributing their enhanced thermal 6 ACS Paragon Plus Environment

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stability as well. In addition, the strong texture and low-angle grain boundaries might also play a role in stabilizing microstructure, and, thus, it would also be interesting to compare thermal and mechanical stability between columnar nc-HEAs and equiaxed nc-HEAs for future studies.

Looking towards to technological applications, these results demonstrate that nanostructured HEAs permit access to high-temperature strength and stability of nanocrystalline metals. As a result, some conventional refractory metals might be replaced by refractory nc-HEAs in future. Their remarkably low strain-rate sensitivity suggest that they are also capable of creep resistance at elevated temperatures. Combined with their high strength-to-density ratios, nc-HEAs are potentially interesting for aerospace, automobile, power, and nuclear industries. Although much work yet remains to optimize nc-HEAs for high-temperature applications, including improving fracture resistance, verifying creep resistance, and testing fatigue behavior, the properties demonstrated here offer strong motivation to further pursue their development and technological usage.

Materials and Methods Bulk polycrystalline (coarse grained) Nb25Mo25Ta25W25 HEAs were prepared from of an equimolar mixture of pure Nb, Mo, Ta and W powders (purity of 99.99%, 99.95%, 99.98% and 99.999%, respectively) using the arc melting technique. Nanocrystalline NbMoTaW HEA films with columnar grains (~70-100 nm grain size) were deposited using the DC magnetron co-sputtering technique. The details are described in Refs. [24, 25], respectively. From the obtained HEA bulk and thin films, the pillar specimens were fabricated using a FIB system (Helios Nanolab 600i, FEI) with a coarse milling condition of 30 kV and 80 pA and a final polishing condition of 5 kV and 24 pA. The FIB-milled pillars (〈011〉 orientation) have a diameter of approximately 1 μm and aspect ratios of about 3. The tapering angle is ~2–4˚, and the top diameters were chosen to calculate engineering stresses.


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The FIB-milled pillars were compressed using a custom-modified in situ high-temperature SEM indenter (Alemnis, Switzerland) in a displacement control mode. Strain-rate jump tests were used for the measurement of strain-rate sensitivity (2×10-4, 5×10-4 10-3, 2×10-3, and 5×10-3). Testing was carried out at 25, 200, 400, and 600 °C. The morphologies of the pillars were characterized using a high resolution SEM (Magellan 400 FEI) after compression. The yield strengths of pillars were measured as offset flow stress at 0.2% of strain. The APT tips of sx- and nc-HEAs were prepared using the FIB system and measured using a LEAP 4000X HR (Cameca) in laser mode with a wavelength of 355 nm, specimen temperature of 40 K, and pulse frequency of 200 kHz.

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Acknowledgements The authors thank S. Gerstl (ScopeM, ETH Zurich) for his help in atom probe analysis, FIRST (ETH Zurich) for sputtering facility, S. Maiti (Laboratory for Crystallography, ETH Zurich) for supplying the bulk HEA sample, Y. Xiao (Laboratory for Nanometallurgy, ETH Zurich) for processing micro-compression data. Y.Z. and H.M. acknowledge financial support through the Swiss National Science Foundation (200021_143633 and 200021_140532). Y. Z. also acknowledge Swiss SNF Early Postdoc.Mobility (P2EZP2_165278).


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Author Contributions Y.Z. prepared the micro-pillar samples, analyzed the data, and prepared the manuscript; J.M.W. carried out the micro-mechanical testing; H.M. sputtered the HEA film; P.O. conducted APT characterization; R.S. supervised the project; all the authors contributed to the discussion and interpretation of the data and the revision of the manuscript.

Competing financial interests The authors declare no competing financial interests

Materials & Correspondence Correspondence should be addressed to [email protected] (Y. Z) or [email protected] (R. S.).

Supporting Information Available: Supplementary figures are available free of charge via the Internet at http://pubs.acs.org.


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Fig. 1. Compression results for the sx- and nc-HEA pillars from room temperature to 600 °C. Representative SEM images of the deformed sx-HEA pillars ((a), (b), (c), and (d)) and nanostructured columnar-grained HEA pillars ((e), (f), (g), and (h)). Corresponding engineering stress–strain curves of (i) the sx-HEA and (j) nc-HEA pillars, showing how flow stresses changes by temperature. Strain-rate jump tests are applied to measure the strain-rate sensitivity using initial and final strain rates of 10-3 s-1 and four other strain rates of 2×10-4 s-1, 2×10-3 s-1, 5×10-4 s-1, and 5×10-3 s-1.


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Fig. 2. Comparison of elevated-temperature strengths in nc-metals. a) Yield strength and (b) the percentage of strength reduction (compared to room-temperature strength) as a function of test temperature for sx- and nc-HEA pillars, indicating that the nc-HEAs exhibit substantially lower strength softening at elevated temperatures than the bulk HEAs, the sx-ones, and nc-Cu-Ta alloys16. (c) Normalized critical resolved shear stress (τ/µ) as a function of homologous temperature (Tt/Tm) for the sx- and nc-HEA pillars in this study, bulk coarse-grained (cg) Nb, Ta, Mo, and W (grain sizes of 100 µm)26, nc-Ni29, nanotwinned (nt) Cu27, ultrafine grained (ufg) Al28, ufg-Cr30, nc-W31, nc-Ta32, nc-Inconel 71833, and nc-Cu-Ta alloys16 tested in the strain rate of ~10-3 s, indicating the nc-HEA pillars exhibit the highest normalized strength (~1/50-1/30) among all the bulk and nanostructured metals (τ is critical resolved shear strength, µ is the corresponding shear modulus, Tt is testing temperature, and Tm is melting temperature; a Schmid factor of 0.5 is used here). (d) The Ashby-inspired map of specific strength (strength-to-density ratio) versus test temperature, showing that the nc-HEAs exhibit the highest strength-to-density ratio in all the nc-metals at the same tested temperatures.


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Fig. 3. Evaluation of strain-rate sensitivity (m) and activation volume (Va) as a function of temperature for sx- and nc-HEAs. (a) Strain-rate sensitivity versus temperature, indicating that the m values of sx-HEAs decrease with increasing temperature, while nc-HEAs are strain-rate insensitive with a consistent m value of ~0.005. (b) Activation volume versus temperature, suggesting that the deformation of nc-HEAs is dominated by kink-pair mechanism in a large temperature range, while sx-HEAs are deformed primarily by the Peierls mechanism at 600 °C.


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Fig. 4. Reconstruction of the APT tips showing elemental distribution in sx- and nc-HEAs. The chemical arrangement maps of (a) a sx-HEA tip and (c), (d) a nc-HEA tip. (b), (e), and (f) are corresponding one-dimensional concentration profiles of (a), (c), and (d), respectively. The APT results indicate a homogenous distribution of principle elements (Nb, Ta, Mo, and W) in both the sx-HEA ((a) and (b)) and the nc-HEA ((c) and (e)). Formation of oxides and nitrides (TaN, TaO, NbN, NbO, and WN, see (d) and (f)) and segregation of foreign elements (N, O, and C, see Supplementary information Fig. S4, ((c-f) are adapted from Ref. [40]) are observed at the grain boundary (GB) region.


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1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

a

sx-25°C

b

nc-25°C

f

sx-200°C

c

sx-400°C

d

sx-600°C

nc-200°C

g

nc-400°C

h

nc-600°C

Nano Letters

sx ~1 µm

~70-100 nm

1 μm

e

nc ~1 µm

j

i

10-3

2×10-4 2×10-3 5×10-4 5×10-3

10-3

ACS Paragon Plus Environment

10-3

2×10-4 2×10-35×10-4 5×10-3 10-3

a 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43

c

b

Nano Letters

Ideal strength

d


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a

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1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19

b

Nano Letters


a 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54

sx-HEA

b

Nano Letters

15 nm

c

e

d nc-HEA GB

15 nm

f


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Page 19 of 19 a

1 2 sx 3 4 5 ~1 µm 1 μm 6 7 8~70-100 nm e 9 10 11 nc 12 13 14~1 µm 15 16 i 17 18 19 20 21 22 23 24 25 26 27 10 2×10 28 29 30 31 -3

Temperature

Nano Letters sx-25°C b

sx-200°C c

sx-400°C d

sx-600°C

nc-25°C f

nc-200°C g

nc-400°C h

nc-600°C GB

j

-4

2×10-3 5×10-4 5×10-3

10-3

10-3


2×10-4 2×10-35×10-4 5×10-3 10-3

Nanocrystalline High-Entropy Alloys: A New Paradigm in High

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