TSpace Research Repository tspace.library.utoronto.ca Nanocrystalline High-Entropy Alloys: A New Paradigm in High- Temperature Strength and Stability Yu Zou, Jeffrey M. Wheeler, Huan Ma, Philipp Okle and Ralph Spolenak Version Post-print/Accepted Manuscript Citation (published version) Zou Y, Wheeler JM, Ma H, Okle P, Spolenak R. Nanocrystalline high- entropy alloys: a new paradigm in high-temperature strength and stability. Nano letters. 2017 Feb 9;17(3):1569-74. Publisher’s Statement The article has been published in final form at [10.1021/acs.nanolett.6b04716] How to cite TSpace items Always cite the published version, so the author(s) will receive recognition through services that track citation counts, e.g. Scopus. If you need to cite the page number of the author manuscript from TSpace because you cannot access the published version, then cite the TSpace version in addition to the published version using the permanent URI (handle) found on the record page. This article was made openly accessible by U of T Faculty. Please tell us how this access benefits you. Your story matters.
15
Embed
Nanocrystalline High-Entropy Alloys: A New Paradigm in ... · Zou Y, Wheeler JM, Ma H, Okle P, Spolenak R. Nanocrystalline high-entropy alloys: a new paradigm in high-temperature
This document is posted to help you gain knowledge. Please leave a comment to let me know what you think about it! Share it to your friends and learn new things together.
Transcript
TSpace Research Repository tspace.library.utoronto.ca
Nanocrystalline High-Entropy Alloys: A New Paradigm in High-
Temperature Strength and Stability
Yu Zou, Jeffrey M. Wheeler, Huan Ma, Philipp Okle and Ralph Spolenak
Version Post-print/Accepted Manuscript
Citation (published version)
Zou Y, Wheeler JM, Ma H, Okle P, Spolenak R. Nanocrystalline high-entropy alloys: a new paradigm in high-temperature strength and stability. Nano letters. 2017 Feb 9;17(3):1569-74.
Publisher’s Statement The article has been published in final form at [10.1021/acs.nanolett.6b04716]
How to cite TSpace items
Always cite the published version, so the author(s) will receive recognition through services that track citation counts, e.g. Scopus. If you need to cite the page number of the author manuscript from TSpace
because you cannot access the published version, then cite the TSpace version in addition to the published version using the permanent URI (handle) found on the record page.
This article was made openly accessible by U of T Faculty. Please tell us how this access benefits you. Your story matters.
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,*
1Laboratory for Nanometallurgy, Department of Materials, ETH Zurich, Vladimir-Prelog-Weg 5, CH-8093 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.).
Inconel 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
5
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 size-
dependent 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 pillars35.
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, 𝜀𝜀̇:
𝑚𝑚 = 𝜕𝜕 𝑙𝑙𝑙𝑙𝜎𝜎𝜕𝜕 𝑙𝑙𝑙𝑙 �̇�𝜀
38. 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 𝑉𝑉𝑎𝑎 =
√3 𝐾𝐾𝐵𝐵𝑇𝑇𝑚𝑚𝜎𝜎
, 38 where kB is the Boltzmann constant of 8.617×10−5 eV/K and T is the absolute temperature. The
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
6
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
7
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.
8
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.
References:
1. Gleiter, H., Nanocrystalline materials. In Advanced Structural and Functional Materials, Springer: 1991; pp 1-37.
2. Meyers, M. A.; Mishra, A.; Benson, D. J. Progress in Materials Science 2006, 51, (4), 427-556. 3. Lu, K.; Lu, L.; Suresh, S. Science 2009, 324, (5925), 349-352. 4. Van Swygenhoven, H.; Weertman, J. R. Mater Today 2006, 9, (5), 24-31. 5. Hall, E. Proceedings of the Physical Society. Section B 1951, 64, (9), 747. 6. Reed, R. C., The superalloys: fundamentals and applications. Cambridge university press: 2008.
1-32 7. Van Swygenhoven, H.; Derlet, P. Phys Rev B 2001, 64, (22), 224105. 8. Shan, Z.; Stach, E. A.; Wiezorek, J. M. K.; Knapp, J. A.; Follstaedt, D. M.; Mao, S. X. Science 2004,
305, (5684), 654-657. 9. Rupert, T.; Gianola, D.; Gan, Y.; Hemker, K. Science 2009, 326, (5960), 1686-1690. 10. Gertsman, V. Y.; Birringer, R. Scripta Metall Mater 1994, 30, (5), 577-581. 11. Mathaudhu, S. N.; Boyce, B. L. Jom-Us 2015, 67, (12), 2785-2787. 12. Liu, X. C.; Zhang, H. W.; Lu, K. Science 2013, 342, (6156), 337-340. 13. Zhang, X.; Misra, A. Scripta Materialia 2012, 66, (11), 860-865. 14. Weissmüller, J. Nanostructured Materials 1993, 3, (1–6), 261-272. 15. Chookajorn, T.; Murdoch, H. A.; Schuh, C. A. Science 2012, 337, (6097), 951-954. 16. Darling, K. A.; Rajagopalan, M.; Komarasamy, M.; Bhatia, M. A.; Hornbuckle, B. C.; Mishra, R. S.;
Solanki, K. N. Nature 2016, 537, (7620), 378-381. 17. Yeh, J. W.; Chen, S. K.; Lin, S. J.; Gan, J. Y.; Chin, T. S.; Shun, T. T.; Tsau, C. H.; Chang, S. Y.
Advanced Engineering Materials 2004, 6, (5), 299-303. 18. Cantor, B.; Chang, I. T. H.; Knight, P.; Vincent, A. J. B. Materials Science and Engineering: A 2004,
375–377, (0), 213-218. 19. Gludovatz, B.; Hohenwarter, A.; Catoor, D.; Chang, E. H.; George, E. P.; Ritchie, R. O. Science
2014, 345, (6201), 1153-1158. 20. Li, Z.; Pradeep, K. G.; Deng, Y.; Raabe, D.; Tasan, C. C. Nature 2016. 21. Dolique, V.; Thomann, A. L.; Brault, P.; Tessier, Y.; Gillon, P. Surface and Coatings Technology
2010, 204, (12–13), 1989-1992.
9
22. Senkov, O. N.; Wilks, G. B.; Scott, J. M.; Miracle, D. B. Intermetallics 2011, 19, (5), 698-706. 23. Giwa, A. M.; Liaw, P. K.; Dahmen, K. A.; Greer, J. R. Extreme Mechanics Letters 2016, 8, 220-228. 24. Zou, Y.; Maiti, S.; Steurer, W.; Spolenak, R. Acta Mater 2014, 65, (0), 85-97. 25. Zou, Y.; Ma, H.; Spolenak, R. Nat Commun 2015, 6. 26. Ashby, M. F., Mechanisms of Deformation and Fracture. In Advances in Applied Mechanics, John,
W. H.; Theodore, Y. W., Eds. Elsevier: 1983; Vol. Volume 23, pp 117-177. 27. Lu, L.; Chen, X.; Huang, X.; Lu, K. Science 2009, 323, (5914), 607-610. 28. Wheeler, J. M.; Maier, V.; Durst, K.; Göken, M.; Michler, J. Materials Science and Engineering: A
2013, 585, 108-113. 29. Wang, Y. M.; Hamza, A. V.; Ma, E. Acta Mater 2006, 54, (10), 2715-2726. 30. Maier, V.; Hohenwarter, A.; Pippan, R.; Kiener, D. Scripta Materialia 2015, 106, 42-45. 31. Wei, Q.; Zhang, H. T.; Schuster, B. E.; Ramesh, K. T.; Valiev, R. Z.; Kecskes, L. J.; Dowding, R. J.;
Magness, L.; Cho, K. Acta Mater 2006, 54, (15), 4079-4089. 32. Wei, Q.; Pan, Z. L.; Wu, X. L.; Schuster, B. E.; Kecskes, L. J.; Valiev, R. Z. Acta Mater 2011, 59, (6),
2423-2436. 33. Sh, M.; Ermachenko, A. Journal of Physics: Conference Series 2010, 240, (1), 012118. 34. Bauer, J.; Schroer, A.; Schwaiger, R.; Kraft, O. Nat Mater 2016, advance online publication. 35. Dimiduk, D. M.; Uchic, M. D.; Parthasarathy, T. A. Acta Mater 2005, 53, (15), 4065-4077. 36. Vitek, V. Progress in Materials Science 1992, 36, (0), 1-27. 37. Mohanty, G.; Wheeler, J. M.; Raghavan, R.; Wehrs, J.; Hasegawa, M.; Mischler, S.; Philippe, L.;
Michler, J. Philos Mag 2014, 1-18. 38. Caillard, D.; Martin, J.-L., Thermally activated mechanisms in crystal plasticity. Elsevier: 2003;
Vol. 8. 39. Wei, Q.; Kecskes, L. J. Materials Science and Engineering: A 2008, 491, (1–2), 62-69. 40. Zou, Y.; Okle, P.; Yu, H.; Sumigawa, T.; Kitamura, T.; Maiti, S.; Steurer, W.; Spolenak, R. Scripta
Materialia 2017, 128, 95-99.
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).
10
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
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.
12
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.
13
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.
14
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.