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A nanocrystalline AlCoCuNi medium-entropy alloy with high thermal stability via entropy and boundary engineering
Materials Science & Engineering A 774 (2020) 138925
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Materials Science & Engineering A journal homepage: http://www.elsevier.com/locate/msea
A nanocrystalline AlCoCuNi medium-entropy alloy with high thermal stability via entropy and boundary engineering H.W. Deng a, b, Z.M. Xie a, **, M.M. Wang a, b, Y. Chen a, b, R. Liu a, J.F. Yang a, T. Zhang c, *, X. P. Wang a, Q.F. Fang a, C.S. Liu a, Y. Xiong d a
Key Laboratory of Materials Physics, Institute of Solid State Physics, Chinese Academy of Sciences, Hefei, Anhui, 230031, China University of Science and Technology of China, Hefei, Anhui, 230026, China School of Physics and Electronic Engineering, Guang Zhou University, Guangzhou, Guangdong, China d State Key Laboratory of Environment-friendly Energy Materials, Southwest University of Science and Technology, Mianyang, 621010, China b c
A R T I C L E I N F O
A B S T R A C T
Keywords: AlCoCuNi Medium-entropy alloy Dual-phase nanocrystalline Thermal stability
A promising strategy involving the synergetic effects of the high-entropy engineering and interface architectures has been proposed to realize the thermal stability of nanocrystalline alloys. A bulk dual-phase nanocrystalline (DPNC-) AlCoCuNi medium-entropy alloy was prepared by mechanical alloying (MA) and spark plasma sintering (SPS) process. These DPNC structures include Cu-rich and AlCoNi-rich nanocrystallines (NCs) with an average grain size of 46 nm. This DPNC-AlCoCuNi material could maintain the nanostructures as well as a high hardness of about 580 HV even after annealing at 900 � C for 50 h. The extremely high thermal stability has been attributed to the extensive thermal-stabled low-energy phase boundaries, low-angle grain boundaries, the high-entropy and sluggish diffusion effects.
1. Introduction Nanocrystalline (NC) alloys, because of the very fine grains and high density of grain boundaries (GBs), exhibit a variety of excellent prop erties such as high strength and hardness, good fatigue resistance and superior radiation tolerance [1–5]. However, the inherent thermal instability of nanograined materials severely not only limits the poten tial application of these materials at elevated temperatures but also makes them hard to process [2,6]. It has been reported that many pure NC metals have a strong spontaneous tendency to grain coarsening, even at relatively low temperature (e.g., room temperature), to release the GB energy of the system [7,8]. Consequently, designing nanocrystalline alloys with desirable thermal stability for elevated temperature appli cations has attracted focused research attention of metallurgists. Recently, Lu et al. [9] proposed a concept of entropy engineering in high-entropy alloys (HEAs), which contain five or more principal ele ments in equal or near equal atomic percent, that could effectively improve the thermal stability of HEAs. This stabilizing way of entropy engineering is essentially thermodynamic stabilization of GBs by reducing the free energy of the system and introducing the solute drag.
From a thermodynamic point of view, the lower GB energy in HEAs may reduce the thermodynamic driving force for coarsening and hence sta bilize the grains [9,10]. In the aspect of kinetic stabilization, considering the presence of and solute-drag effect of many different adsorbates, coarsening kinetics can be suppressed [11]. Moreover, the thermody namic driving force for grain growth and precipitation can be reduced the sluggish diffusion kinetics in HEAs [12]. Furthermore, another approach for stabilizing nanostructure alloys by modifying the architectures or introducing low-energy interfaces, such as low-angle GBs, twin boundaries, coherent phase boundaries (PBs) and immiscible PBs, is proposed [4]. For example, nanostructures with a layered arrangement of low-angle coherent nanotwin boundaries in sputtered Cu films exhibited better thermal stability than monolithic nanocrystals with high-angle GBs [13]. Recently, it has been reported that nanostructured multi-phase materials like Cu/Ta, Cu/Nb nano-lamellar multilayers, which contain bi- and/or tri-phase interfaces comprised of immiscible metals, demonstrate excellent thermal stabil ities due to the lower PB energies than that of the GB [14–16]. Hence, a promising strategy involving the synergetic effects of the high-entropy engineering and interface architectures would prospectively realize a
* Corresponding author. ** Corresponding author. E-mail addresses: [email protected] (Z.M. Xie), [email protected] (T. Zhang). https://doi.org/10.1016/j.msea.2020.138925 Received 13 November 2019; Received in revised form 4 January 2020; Accepted 4 January 2020 Available online 7 January 2020 0921-5093/© 2020 Elsevier B.V. All rights reserved.
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high thermal stability of NC alloys. To establish the validity of this strategy, in this work, a bulk AlCo CuNi medium-entropy nanocrystalline alloy consist of a dual-phase (Curich FCC and AlCoNi-rich BCC) structure has been designed and fabri cated. To avoid the grain coarsening and precipitation formation in traditional casting, here this dual-phase nanocrystalline (DPNC-) medium-entropy alloy (MEA) was manufactured by mechanical alloying (MA) followed by consolidation via spark plasma sintering (SPS). The alloy design strategy and detailed fabrication processes are presented. The thermal stability of this DPNC-MEA was investigated and the cor responding mechanism has been proposed.
ball-milling machine. Al, Co, Cu, Ni powders with an equal-atomic ratio (300 mesh, purity > 99.9%, purchased from Aladdin Reagent Co., LTD) and tungsten carbide balls were placed in stainless-steel vials with a powder-to-ball weight ratio of 1:10. Milling was carried out for 50 h in an argon atmosphere with a rotation speed of 300 rpm with a two-step method: firstly, a dry milling for 40 h without a process control agent (PCA) and then wet milling for 10 h with ethanol as a PCA. Subse quently, the mechanical-alloyed AlCoCuNi powder was consolidated by SPS (furnace SE-607, FCT Group, Germany) at 900 � C, with a heating � rate of 100 C/min and a constant pressure of 50 MPa. 2.3. Mechanical and microstructure characterization
2. Material and methods
The DPNC-AlCoCuNi specimens were annealed at a temperature ranging from 500 to 1100 � C for various time to investigate the thermal stability. Then the annealed samples were subjected to Vickers micro hardness test at room temperature with a load of 500 g and a dwell time of 15 s. Each examined sample was indented for at least 8 times in different areas. The compressive tests were carried out using an Instron5967 machine using a strain rate of 1 � 10 3 s 1. The phase composition was determined by X-ray diffraction (XRD, PANalytical X’Pert Pro MRD diffractometer) by using Cu Kα radiations (wavelength λ ¼ 0.154 nm) at a voltage of 45 kV and a tube current of 200 mA. The microstructural characterization was carried out using scanning electron microscope (SEM, SU8020) and transmission electron microscope (TEM, Tecnai G2 F20) with an energy-dispersive X-ray spectrometer (EDS). Thin speci mens for TEM observation were firstly dimpled and then polished using a Gatan 691 precision ion-milling machine with a low energy ion-beam (3 keV).
2.1. The alloy design strategy Since the concept of high-entropy alloys were proposed in 2004 [17], most of the related studies have mainly focused on the design of single-phase high-entropy alloy (HEA) components via maximized configurational entropy to stabilize single-phase structures. Recent re sults show that a high-entropy dual-phase (DP-) HEAs with a high degree of coherent interfaces between two phases exhibit an improved strength as well as excellent ductility, especially at high temperatures, compared to the single-phase HEAs [18]. More importantly, the DP-structure will suppress the recrystallization and retard the grain coarsening even at high temperatures that could improve the thermal stability of materials [19]. Hence, it is promising to explore new DP-HEAs system for the application as high-temperature structural materials. It has been reported that there are several DP-HEA components, such as AlCoCrCu0.5Fe, AlxCoCrFeNi and Al0.5CoCrCuFeNi [20–22]. In these systems with high Fe content, DP-structure is unstable at high temper atures. However, the DP-structure in a Fe-free AlCoCrCuNi alloy can be maintained up to 1200 � C [23]. It is known that the Cr element has a strong tendency to form a brittle σ phase or segregate at grain bound aries during the fabrication process. So in this work, AlCoCuNi medium-entropy alloy (MEA) with Cr element free was designed to avoid the formation of precipitate phases. Besides, to ensure the for mation of stabilized solid-solution phases without additional bulk-metallic glass, empirical conditions of Ω � 1.1 and δ � 6.6% should be satisfied, where Ω is an indicator of the stability of the multicom ponent solid solution and δ is the atomic size difference of the compo nent elements [24], rffiffiffiffiffiffiffi�ffiffiffiffiffiffiffiffiffiffiffiffiffiffi�ffiffiffiffi ri 2 δ ¼ Σci 1 r
3. Results and discussion 3.1. Microstructure and chemical composition of 50 h ball milling powder The SEM image of the as-milled powder and corresponding EDS mapping are shown in Fig. 1. The raw Al, Co, Cu, Ni powder was crushed down and repeatedly welded together to form a uniform distribution of elemental constituent during the first 40 h dry milling. And then a period of 10 h wet milling with ethanol as a PCA is employed to avoid the aggregation of the powders and produce finer particles. After the 50 h milling, powder particles exhibit near-equiaxed and irregular morphology, with particle sizes varying from 5 to 25 μm, as shown in Fig. 1a. Furthermore, the corresponding EDS mapping reveals that each milled powder particle contains homogeneous distribution of Al, Co, Cu and Ni elements as shown in Fig. 1b.
where r ¼ Σci ri is the average atomic radius, ci, ri is the molar content and atomic radius of the ith component, respectively. Tm ¼ Σci ðTm Þi is the average melting point of the alloy system. ΔSmix ¼ RΣci lnci is the is the entropy of mixing for equimolar ratio alloy. ΔH ¼ Σi
3.2. Phase evolution of the AlCoCuNi MEA after MA, SPS, and annealing The manually grinded AlCoCuNi mixture powder with an equalatomic ratio was characterized by X-ray diffraction (XRD) to exhibit Al, Co, Cu, and Ni phases, as shown in Fig. 2a “0 h” sample. After high energy ball milling, diffraction peaks of Cu-based FCC appear and peaks of Co, Cu and Ni are absent, as shown in Fig. 2a “50 h milled” sample. In addition, a new series of diffraction peaks were detected and indexed to the AlCo-based BCC structure, as indicated by red arrows in Fig. 2a. The alloy formation is governed by the inherent thermodynamic and kinetic constraints of the system like binary enthalpy of mixing (ΔHmix) [26]. Phase separation is often observed in HEAs when a high entropy of mixing cannot counteract a significantly positive enthalpy of mixing (or formation of significantly negative enthalpy induced by intermetallic phases). It is suggested that during MA, Co and Ni have a strong ten dency to alloying with Al and produce an AlCoNi-rich BCC phase, owing to the high negative ΔHmix (Fig. 2b), which confirms the XRD results. Also, all peaks of Cu-based FCC and AlCo-based BCC are obviously broadened after ball milling, suggesting a remarkable grain refinement and lattice distortion. These results imply that the formation of
enthalpy of mixing. To construct a thermal stabled dual-phase MEA alloy, the chemical composition of Al25Co25Cu25Ni25 with Ω ¼ 2.22 and δ ¼ 5.82% (using data in Table 1 and Fig. 2b) was designed according to this concept [25]. 2.2. Fabrication methods DPNC-AlCoCuNi MEA powder was synthesized by mechanical alloying (MA) from pure elemental powders in a high-energy planetary
Table 1 Atomic radius, melting temperature of Al, Co, Cu, Ni pure elements. r (nm) Tm (K)
Al
Co
Cu
Ni
0.14317 933
0.12510 1768
0.1278 1358
0.12459 1728
2
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Materials Science & Engineering A 774 (2020) 138925
Fig. 1. (a) SEM image of 50 h as-milled AlCoCuNi powders, and (b) corresponding EDS mapping of Al, Co, Cu, Ni, exhibiting the homogeneous distribution.
Fig. 2. (a) XRD patterns of AlCoCuNi mixture powders before and after high energy ball milling, and SPSed bulk DPNC-AlCoCuNi MEA, (b) the value of enthalpy of mixing (ΔHmix) for atomic pairs [32], (c) the (111) diffraction peak of Cu-rich FCC in bulk DPNC-AlCoCuNi and the relative standard (111) diffraction peak in pure Cu.
Fig. 3. TEM characterizations of (a)–(c) the as SPSed and (d)–(e) 900 � C annealed bulk DPNC-AlCoCuNi MEA: (a, d) bright-field images, (b, e) dark-field images and (f) grain size distributions. 3
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DPNC-AlCoCuNi MEA powder after MA. An effective SPS technique was employed to further consolidate this powder and to get a more equi librium bulk DPNC-AlCoCuNi. The corresponding XRD result of the SPSed bulk sample revealed a more obvious dual-phase structure with Cu-rich FCC (lattice constant, 0.36081 nm) and AlCoNi-rich BCC (lattice constant, 0.28646 nm) phases. The peak position of the FCC phase slightly shifts to the low angle in comparison with pure Cu (Fig. 2c), indicating partial dissolution of larger sized atoms (Al, Table 1) resulting in expansion of FCC lattice. The bright-field and dark-field TEM images of the SPSed AlCoCuNi MEA are presented in Fig. 3a and Fig. 3b, respectively. The SPSed specimens show a NC structure with a mean grain size of about 46 nm. It is worth mentioning that the grain size in both Cu-rich FCC and AlCoNirich BCC was counted as a single species because of their comparable crystalline sizes and distributions. Meanwhile, from the corresponding selected area electron diffraction (SAED) pattern (Fig. 3c), it can be seen that some diffraction rings belong to both FCC and BCC phase which suggest the existence of similar interplanar crystal spaces and forebodes ideal crystallographic orientations between these two phases. After relatively high-temperature annealing at 900 � C for 1 h, the NC structure was maintained and the mean grain size kept at 52 nm, suggesting un precedented thermal stability in our DPNC-AlCoCuNi MEA. The com parable histograms of the grain size distribution (Fig. 3f) in the as SPSed and annealed ones further confirm the high structural stability. The grain size distribution is calculated from more than 300 grains from several TEM images captured from different regions of the samples.
critical temperature of about 900 � C, which is significantly higher than that of other NC materials as indicated by the open blue arrows in Fig. 4b. To further evaluate the thermal stability of DPNC-AlCoCuNi MEA, a series of long-duration annealing up to 50 h was performed, the Vickers hardness dependence of annealing time at various temperatures are shown in Fig. 5. In the case of 900 � C annealing with increasing the annealing time to 10 h, there is just a slight decrease of the harness to 591 HV (ΔHV ~ 8.5%, from the initial 645 HV). It is worth noting that the change in hardness value is negligible for specimens annealed at 900 � C when further increasing the annealing time to 20 h or even to 50 h. For instance, after annealing at 900 � C for 50 h, the hardness of the annealed specimen is still as high as 580 HV. The SEM image of the related fracture surface reveals that the nanostructure of the alloy is still maintained after annealing at 900 � C for 50 h, except for a few grown micron-sized grains as shown in Fig. 5. The newly grown micron-sized grains are Cu-rich FCC structures as illustrated in the following TEM analysis. In the case of 950 � C annealing, a moderate decrease of the harness to 548 HV (ΔHV ~ 15%) is observed after annealing for 10 h. Further increasing the annealing (at 950 � C) time to 50 h, a high level of hardness value of about 508 HV is maintained. The corresponding microstructure of the fracture surface shows that most of the nano structure is retained even after annealing at 950 � C for 50 h, except for some grown micron-sized grains as shown in Fig. 5. While after 10 h annealing at 1000 � C, the hardness is decreased down to 480 HV, about 26% lower than that of the SPSed one, suggesting a typical micro structure instability as indicated in fracture surface morphology shown in the inset of Fig. 5. All these indicate the DPNC-AlCoCuNi MEA indeed possesses an excellent resistance to long-term annealing softening when the annealing temperature is below 900 � C.
3.3. Mechanical properties and stabilities of the DPNC-AlCoCuNi MEA The engineering stress-strain curve of DPNC-AlCoCuNi MEA under uniaxial compression at room temperature is presented in Fig. 4a. The specimen shows an extremely high compressive strength of 2.0 GPa with the strain up to 10.1%, which is higher than that of reported AlCuCr FeMnW0.5(1510 MPa) [27], AlCuNiFeCr (1960 MPa) [28], and Al0.3FeNiCo1.2CrCu (1635 MPa) [29]. It is well known that the hardness is closely related to the microstructure, therefore the structural stability can be further reflected by the evolution of hardness with annealing temperatures. Our DPNC-AlCoCuNi MEA exhibits extremely higher hardness up to 645 HV than that of other NC materials like the NC-CoCrFeMnNi HEA, NC-CoCrFeNi MEA and some Cu-based NC metals and alloys like NC-Cu and Cu/Ta nano-lamellar multilayer, as compared in Fig. 4b. More importantly, the DPNC-AlCoCuNi MEA maintains this high hardness value even after 900 � C annealing for 1 h, suggesting the outstanding thermal stability. Further increasing the annealing tem perature to 1000 � C, the hardness just starts decreasing to 586 HV because of the structure failure of the nanosized Cu-rich FCC. That’s to say, the DPNC-AlCoCuNi MEA could maintain an unprecedented ther mal stabled nano-structure as well as a high strength/hardness until a
Fig. 5. Vickers hardness evolution of DPNC-AlCoCuNi MEA after annealing at 900, 950, and 1000 � C for the time rang of 1–50 h; the SEM images and the inset show the nanostructure of the alloy annealing at 950 � C for 50 h.
Fig. 4. (a) Engineering stress-strain curve of the bulk SPSed DPNC-AlCoCuNi MEA specimen under uniaxial compression at room temperature. (b) Vickers microhardness evolution of DPNC-AlCoCuNi MEA after annealing for 1 h in the temperature range of 500–1100 � C, values of other NC-materials are present for comparison. 4
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Fig. 6 compares the SEM microstructure of the fracture surface after annealing at 900, 950, and 1000 � C for 5 and 10 h. It can be seen that the grain structure consists of nanostructured grains and recrystallized coarse grains (as indicated by the black open arrows in Fig. 6). The proportion of recrystallization under various annealing conditions was quantified using an image recognition algorithm and the results are presented in Fig. 6. A negligible recrystallization proportions of about 3.5% and 4.3% are observed after annealing at 900 � C for 5 h and 10 h, respectively. And the proportion of recrystallization increases with increasing the annealing temperature from 900 � C to 1000 � C, which is consistent with the decreasing trend of hardness. For example, recrys tallization proportion increased to 23.3% after annealing at 1000 � C for 10 h, which is 6 times than of the 900 � C annealed one.
profile confirms this dual-phase nanocrystalline separation structure, as depicted in Fig. 7b. This DPNC separation structure was bred from MA and consolidated during SPS, resulting in low-energy states in both Curich and AlCoNi-rich NCs because of the high entropy effects. In addi tion, these extensively separated PBs could generate a sufficient solutedrag effect, suppress kinetics for coarsening and hence significantly in crease the thermal stabilities. The HRTEM image (Fig. 7c) provides further evidence that Cu-rich and AlCoNi-rich NCs constructed a continuous PB as indicated by the white dotted line and the inset cor responding fast Fourier transform (FFT) images. The representative in verse fast Fourier transform (IFFT) of the yellow dotted square region in Fig. 7c is presented in Fig. 7d to highlight the details of the PB. An evidently special orientation for the FCC and BCC crystals with (31–1)FCC//(01-1)BCC and [011]FCC//[-111]BCC is observed. The present case shows that the orientation is close to the Kurdjumov-Sachs (K–S) relationships [30] but contains an 18� misorientation between the close-packed planes of (1–11)FCC and (110)BCC. In addition, the results calculated from diffraction spots of the PB region exhibit a negligible lattice mismatch (δ ¼ 0.4%) between the parallel planes of (31–1)FCC and (02-2)BCC, suggesting the formation of a coherent low-energy PB. It is interesting to mention that the lattice mismatch between the (31–1)FCC and (02-2)BCC is calculated to be δ~6.9% from the matrix region (away from the PB, as shown in Fig. 7c inset FFT images), which is quite larger than the case of PB region and beyond the coherent scope. However, a slight decrease of lattice space of (31–1)FCC and an increase
3.4. Interface structure of the DPNC-AlCoCuNi MEA To further comprehend essential factors leading to the high thermal stability in this DPNC-AlCoCuNi MEA, extensive interfaces between Curich FCC and AlCoNi-rich BCC phases were detailedly investigated by high-angle annular dark-field scanning transmission electron micro scopy (HAADF-STEM), high-resolution TEM (HRTEM) and EDS. The HAADF-STEM image (Fig. 7a) accentuates the Cu-rich nanograins on the basis of atomic-number contrast, which were separated and surrounded by AlCoNi-rich nanograins in the SPSed DPNC-AlCoCuNi MEA. A locally amplified image with the corresponding HAADF-STEM-EDS line-scan
Fig. 6. SEM microstructures of DPNC-AlCoCuNi MEA after annealing at 900, 950, and 1000 � C for the 5 and 10 h; and the proportion of recrystallization are presented. 5
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Fig. 7. Microstructure of DPNC-AlCoCuNi MEA: (a) a HAADF image, (b) HAADF-STEM-EDS line-scan across three different nanograins, (c) HRTEM image showing the PB, (d) IFFT from the yellow dotted square region in (c). (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
of (02-2)BCC have been triggered through a self-adaptation process to decrease the interface energy and thus a stabled low-energy coherent- state PB has been constructed. This self-adaptation process has been indicated from the atomic redistribution of partial Al atoms (with the largest atomic radius) [31], which diffuse from the Cu-rich FCC to AlCoNi-rich BCC during high-temperature consolidation to reduce the PB interface energy. This Al diffusion could be confirmed by the corre sponding HADDF-STEM-EDS line-scan profile of Al distribution (the blue line in Fig. 7b) which has a gentle slope through the Cu-rich FCC to AlCoNi-rich BCC interface. In addition to these extensive coherent PBs, a few of other lowenergy interfaces like low-angle GBs in Cu-rich FCC were also observed, as presented in Fig. 8. Two Cu-rich FCC grains were viewed along the z ¼ [011]FCC, as indicated by the inset corresponding FFT images. The corresponding HRTEM image (Fig. 8a) shows a low-angle GB in the Cu-rich FCC phase as indicated by the white dotted line. In this case, some lattice distortion, dislocations, and a resulting disorien tation of about 8� were observed, as presented in Fig. 8b. These lowangle GBs with low energy states in Cu-based alloys have been indi cated exhibiting better thermal stability than monolithic nanocrystals with high-angle grain boundaries [29]. Combining the XRD analysis, the chemical composition, and the microstructure characterization, the formation process of this thermally stabled DPNC-AlCoCuNi MEA was summarized, as illustrated schemat ically in Fig. 9. After 50 h ball milling, Al, Co, Cu and Ni pure elements with an equal-atomic ratio roughly constructed the DP structure (phase separation) with both the AlCoNi-rich BCC and Cu-rich FCC phases, as indicated by the blue lines in Fig. 9b. In addition, remarkable grain re finements to form nanocrystalline grains/cells (indicated by the black dotted lines in Fig. 9b) could be confirmed by the obviously broadened XRD diffraction peaks as shown in Fig. 2a. Finally, these extensive nanosized PBs between the Cu-rich FCC to AlCoNi-rich BCC phases and some nano-sized low-angle GBs were preserved and further consolidated
(reducing the lattice defects like vacancies and dislocations) to form the stabled bulk DPNC-AlCoCuNi MEA during the effective SPS process, as shown in Fig. 9c. Hence, it is comprehensible that the bulk DPNCAlCoCuNi MEA achieves the unprecedented thermal stability: i) MA constructing the DPNC MEA powder with Cu-rich FCC and AlCoNi-rich BCC NCs with both the high entropy and sluggish diffusion effects, ii) high-temperature SPS consolidating DPNC structure through a selfadaptation process, resulting in extensive thermal-stabled low-energy coherent PBs and some low-angle GBs. As a representative example, the microstructure after 1000 � C annealing for 1 h is shown in Fig. 10. The DPNC-AlCoCuNi MEA is hard to maintain the nanostructure due to the instability of Cu-rich FCC. Correspondingly, the Cu-rich FCC nanocrystalline grains begin to grow into micron ones as highlighted by the white dotted line in Fig. 10. It is worth mentioning that AlCoNi-rich BCC nanocrystalline grains are still retained even after the 1000 � C annealing. 4. Conclusions A bulk DPNC-AlCoCuNi MEA with unprecedented thermal stability was successfully processed via MA and SPS consolidation. This NC MEA has an average grain size of 46 nm, an extremely high compressive strength of 2.0 GPa with the strain up to 10.1% and a high Vickers microhardness of 645 HV. More importantly, this DPNC-AlCoCuNi material could maintain the nanostructures as well as a high hardness of about 580 HV even after annealing at 900 � C for 50 h. This critical temperature of about 900 � C is significantly higher than that of other Cu containing NC materials and NC MEA alloys [15, 16, 27, 31–34]. The extremely high thermal stability has been indicated from i) the thermalstabled matrix including Cu-rich FCC and AlCoNi-rich BCC NCs with high entropy and sluggish diffusion effects and ii) extensive thermalstabled low-energy coherent PBs and some low-angle GBs.
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Fig. 8. (a) High-resolution TEM (HR-TEM) of the grain boundary, (b) Inverse fast Fourier transform (IFFT) image from the region indicated by the yellow dotted square in (a). The fast Fourier transform (FFT) images in the inset of (a) show the two FCC grains. The white dotted line indicates the grain boundary. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
Fig. 9. A Schematic diagram showing the stable DPNC-AlCoCuNi MEA formation process.
Acknowledgements This work was financially supported by the National Natural Science Foundation of China (Grant Nos.: 51771184, 11735015, 11575241, 51801203, 11674319), the Natural Science Foundation of Anhui Prov ince (Grant No. 1808085QE132) and the Open Project of State Key Laboratory of Environment friendly Energy Materials (Grant No. 18KFHG02). References [1] C.A. Schuh, T.G. Nieh, H. Iwasaki, The effect of solid solution W additions on the mechanical properties of nanocrystalline Ni, Acta Mater. 51 (2) (2003) 431–443. [2] K.S. Kumar, H. Van Swygenhoven, S. Suresh, Mechanical behavior of nanocrystalline metals and alloys11The golden jubilee issue—selected topics in materials science and engineering: past, present and future, in: S. Suresh (Ed.), Acta Mater. 51 (19) (2003) 5743–5774. [3] Y.F. Shen, L. Lu, Q.H. Lu, Z.H. Jin, K. Lu, Tensile properties of copper with nanoscale twins, Scr. Mater. 52 (10) (2005) 989–994. [4] K. Lu, Stabilizing nanostructures in metals using grain and twin boundary architectures, Nat. Rev. Mater. 1 (5) (2016). [5] L.F. Zeng, P. Fan, L.F. Zhang, R. Gao, Z.M. Xie, Q.F. Fang, X.P. Wang, D.Q. Yuan, T. Zhang, C.S. Liu, He irradiation effects in bulk Cu/V nanolayered composites fabricated by cross accumulative roll bonding, J. Nucl. Mater. 508 (2018) 354–360. [6] H.J. Fecht, Intrinsic instability and entropy stabilization of grain boundaries, Phys. Rev. Lett. 65 (5) (1990) 610. [7] B. Günther, A. Kumpmann, H.D. Kunze, Secondary recrystallization effects in nanostructured elemental metals, Scr. Metall. Mater. 27 (7) (1992) 833–838. [8] C.C. Koch, R.O. Scattergood, K.A. Darling, J.E. Semones, Stabilization of nanocrystalline grain sizes by solute additions, J. Mater. Sci. 43 (23–24) (2008) 7264–7272. [9] Z. Lei, X. Liu, H. Wang, Y. Wu, S. Jiang, Z. Lu, Development of advanced materials via entropy engineering, Scr. Mater. 165 (2019) 164–169. [10] Y. Zou, H. Ma, R. Spolenak, Ultrastrong ductile and stable high-entropy alloys at small scales, Nat. Commun. 6 (2015) 7748. [11] K.Y. Tsai, M.H. Tsai, J.W. Yeh, Sluggish diffusion in Co–Cr–Fe–Mn–Ni high-entropy alloys, Acta Mater. 61 (13) (2013) 4887–4897. [12] N. Zhou, T. Hu, J. Huang, J. Luo, Stabilization of nanocrystalline alloys at high temperatures via utilizing high-entropy grain boundary complexions, Scr. Mater. 124 (2016) 160–163. [13] O. Anderoglu, A. Misra, H. Wang, X. Zhang, Thermal stability of sputtered Cu films with nanoscale growth twins, J. Appl. Phys. 103 (9) (2008), 094322.
Fig. 10. A TEM image showing the 1000 � C annealed DPNC-AlCoCuNi MEA.
Declaration of competing interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. CRediT authorship contribution statement H.W. Deng: Conceptualization, Methodology, Investigation, Data curation, Writing - original draft. Z.M. Xie: Methodology, Formal analysis, Data curation, Writing - review & editing. M.M. Wang: Investigation. Y. Chen: Validation. R. Liu: Resources, Supervision. J.F. Yang: Formal analysis, Data curation. T. Zhang: Project administration, Funding acquisition, Writing - review & editing. X.P. Wang: Supervi sion. Q.F. Fang: Project administration, Funding acquisition. C.S. Liu: Project administration, Funding acquisition. Y. Xiong: Funding acquisition.
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Materials Science & Engineering A 774 (2020) 138925 [24] Y. Zhang, T.T. Zuo, Z. Tang, M.C. Gao, K.A. Dahmen, P.K. Liaw, Z.P. Lu, Microstructures and properties of high-entropy alloys, Prog. Mater. Sci. 61 (2014) 1–93. [25] Z. Fu, L. Jiang, J.L. Wardini, B.E. MacDonald, H. Wen, W. Xiong, D. Zhang, Y. Zhou, T.J. Rupert, W. Chen, E.J. Lavernia, A high-entropy alloy with hierarchical nanoprecipitates and ultrahigh strength, Sci. Adv. 4 (10) (2018), eaat8712. [26] M. Vaidya, A. Prasad, A. Parakh, B.S. Murty, Influence of sequence of elemental addition on phase evolution in nanocrystalline AlCoCrFeNi: novel approach to alloy synthesis using mechanical alloying, Mater. Des. 126 (2017) 37–46. [27] D. Kumar, V.K. Sharma, Y.V.S.S. Prasad, V. Kumar, Materials-structure-property correlation study of spark plasma sintered AlCuCrFeMnWx (x ¼ 0, 0.05, 0.1, 0.5) high-entropy alloys, J. Mater. Res. 34 (5) (2019) 767–776. [28] A.I. Yurkova, V.V. Cherniavsky, V. Bolbut, M. Krüger, I. Bogomol, Structure formation and mechanical properties of the high-entropy AlCuNiFeCr alloy prepared by mechanical alloying and spark plasma sintering, J. Alloy. Comp. 786 (2019) 139–148. [29] S. Yang, J. Pi, W. Yang, H. Zhou, D. Zhu, Deformation twinning structure and interface in a FCC-based Al0.3FeNiCo1.2CrCu high-entropy alloy matrix composites, Mater. Lett. 214 (2018) 50–52. [30] J.H. van der Merwe, G.J. Shiflet, The role of structural ledges at phase boundaries—III. F.C.C.-B.C.C. interfaces in Kurdjumov-Sachs orientation, Acta Metall. Mater. 42 (4) (1994) 1199–1205. [31] L.J. Santodonato, Y. Zhang, M. Feygenson, C.M. Parish, M.C. Gao, R.J.K. Weber, J. C. Neuefeind, Z. Tang, P.K. Liaw, Deviation from high-entropy configurations in the atomic distributions of a multi-principal-element alloy, Nat. Commun. 6 (2015) 5964. [32] A. Takeuchi, A. Inoue, Classification of bulk metallic glasses by atomic size difference, heat of mixing and period of constituent elements and its application to characterization of the main alloying element, Mater. Trans. 46 (12) (2005) 2817–2829.
[14] Y. Chen, Y. Liu, F. Khatkhatay, C. Sun, H. Wang, X. Zhang, Significant enhancement in the thermal stability of nanocrystalline metals via immiscible triphases, Scr. Mater. 67 (2) (2012) 177–180. [15] E. Botcharova, J. Freudenberger, L. Schultz, Mechanical and electrical properties of mechanically alloyed nanocrystalline Cu–Nb alloys, Acta Mater. 54 (12) (2006) 3333–3341. [16] L.F. Zeng, R. Gao, Q.F. Fang, X.P. Wang, Z.M. Xie, S. Miao, T. Hao, T. Zhang, High strength and thermal stability of bulk Cu/Ta nano lamellar multilayers fabricated by cross accumulative roll bonding, Acta Mater. 110 (2016) 341–351. [17] J.W. Yeh, S.K. Chen, S.J. Lin, J.Y. Gan, T.S. Chin, T.T. Shun, C.H. Tsau, S.Y. Chang, Nanostructured high-entropy alloys with multiple principal elements: novel alloy design concepts and outcomes, Adv. Eng. Mater. 6 (5) (2004) 299–303. [18] K.R. Lim, K.S. Lee, J.S. Lee, J.Y. Kim, H.J. Chang, Y.S. Na, Dual-phase high-entropy alloys for high-temperature structural applications, J. Alloy. Comp. 728 (2017) 1235–1238. [19] J. Hou, X. Shi, J. Qiao, Y. Zhang, P.K. Liaw, Y. Wu, Ultrafine-grained dual phase Al0.45CoCrFeNi high-entropy alloys, Mater. Des. 180 (2019) 107910. [20] W.-R. Wang, W.-L. Wang, S.-C. Wang, Y.-C. Tsai, C.-H. Lai, J.-W. Yeh, Effects of Al addition on the microstructure and mechanical property of AlxCoCrFeNi highentropy alloys, Intermetallics 26 (2012) 44–51. [21] C.W. Tsai, M.H. Tsai, J.W. Yeh, C.C. Yang, Effect of temperature on mechanical properties of Al0.5CoCrCuFeNi wrought alloy, J. Alloy. Comp. 490 (1) (2010) 160–165. [22] K.E. Guan-Yu, Swe-Kai CHEN, Tung HSU, Jien-Wei YEH, FCC and BCC equivalents in as-cast solid solutions of Al x Co y Cr z Cu 0.5 Fe v Ni w high-entropy alloys, Proc. Prehist. Soc. 31 (2) (2006), 633-381. [23] Ł. Rogal, On the microstructure and mechanical properties of the AlCoCrCuNi high entropy alloy processed in the semi-solid state, Mater. Sci. Eng. A 709 (2018) 139–151.
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Contents lists available at ScienceDirect
Materials Science & Engineering A journal homepage: http://www.elsevier.com/locate/msea
A nanocrystalline AlCoCuNi medium-entropy alloy with high thermal stability via entropy and boundary engineering H.W. Deng a, b, Z.M. Xie a, **, M.M. Wang a, b, Y. Chen a, b, R. Liu a, J.F. Yang a, T. Zhang c, *, X. P. Wang a, Q.F. Fang a, C.S. Liu a, Y. Xiong d a
Key Laboratory of Materials Physics, Institute of Solid State Physics, Chinese Academy of Sciences, Hefei, Anhui, 230031, China University of Science and Technology of China, Hefei, Anhui, 230026, China School of Physics and Electronic Engineering, Guang Zhou University, Guangzhou, Guangdong, China d State Key Laboratory of Environment-friendly Energy Materials, Southwest University of Science and Technology, Mianyang, 621010, China b c
A R T I C L E I N F O
A B S T R A C T
Keywords: AlCoCuNi Medium-entropy alloy Dual-phase nanocrystalline Thermal stability
A promising strategy involving the synergetic effects of the high-entropy engineering and interface architectures has been proposed to realize the thermal stability of nanocrystalline alloys. A bulk dual-phase nanocrystalline (DPNC-) AlCoCuNi medium-entropy alloy was prepared by mechanical alloying (MA) and spark plasma sintering (SPS) process. These DPNC structures include Cu-rich and AlCoNi-rich nanocrystallines (NCs) with an average grain size of 46 nm. This DPNC-AlCoCuNi material could maintain the nanostructures as well as a high hardness of about 580 HV even after annealing at 900 � C for 50 h. The extremely high thermal stability has been attributed to the extensive thermal-stabled low-energy phase boundaries, low-angle grain boundaries, the high-entropy and sluggish diffusion effects.
1. Introduction Nanocrystalline (NC) alloys, because of the very fine grains and high density of grain boundaries (GBs), exhibit a variety of excellent prop erties such as high strength and hardness, good fatigue resistance and superior radiation tolerance [1–5]. However, the inherent thermal instability of nanograined materials severely not only limits the poten tial application of these materials at elevated temperatures but also makes them hard to process [2,6]. It has been reported that many pure NC metals have a strong spontaneous tendency to grain coarsening, even at relatively low temperature (e.g., room temperature), to release the GB energy of the system [7,8]. Consequently, designing nanocrystalline alloys with desirable thermal stability for elevated temperature appli cations has attracted focused research attention of metallurgists. Recently, Lu et al. [9] proposed a concept of entropy engineering in high-entropy alloys (HEAs), which contain five or more principal ele ments in equal or near equal atomic percent, that could effectively improve the thermal stability of HEAs. This stabilizing way of entropy engineering is essentially thermodynamic stabilization of GBs by reducing the free energy of the system and introducing the solute drag.
From a thermodynamic point of view, the lower GB energy in HEAs may reduce the thermodynamic driving force for coarsening and hence sta bilize the grains [9,10]. In the aspect of kinetic stabilization, considering the presence of and solute-drag effect of many different adsorbates, coarsening kinetics can be suppressed [11]. Moreover, the thermody namic driving force for grain growth and precipitation can be reduced the sluggish diffusion kinetics in HEAs [12]. Furthermore, another approach for stabilizing nanostructure alloys by modifying the architectures or introducing low-energy interfaces, such as low-angle GBs, twin boundaries, coherent phase boundaries (PBs) and immiscible PBs, is proposed [4]. For example, nanostructures with a layered arrangement of low-angle coherent nanotwin boundaries in sputtered Cu films exhibited better thermal stability than monolithic nanocrystals with high-angle GBs [13]. Recently, it has been reported that nanostructured multi-phase materials like Cu/Ta, Cu/Nb nano-lamellar multilayers, which contain bi- and/or tri-phase interfaces comprised of immiscible metals, demonstrate excellent thermal stabil ities due to the lower PB energies than that of the GB [14–16]. Hence, a promising strategy involving the synergetic effects of the high-entropy engineering and interface architectures would prospectively realize a
* Corresponding author. ** Corresponding author. E-mail addresses: [email protected] (Z.M. Xie), [email protected] (T. Zhang). https://doi.org/10.1016/j.msea.2020.138925 Received 13 November 2019; Received in revised form 4 January 2020; Accepted 4 January 2020 Available online 7 January 2020 0921-5093/© 2020 Elsevier B.V. All rights reserved.
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high thermal stability of NC alloys. To establish the validity of this strategy, in this work, a bulk AlCo CuNi medium-entropy nanocrystalline alloy consist of a dual-phase (Curich FCC and AlCoNi-rich BCC) structure has been designed and fabri cated. To avoid the grain coarsening and precipitation formation in traditional casting, here this dual-phase nanocrystalline (DPNC-) medium-entropy alloy (MEA) was manufactured by mechanical alloying (MA) followed by consolidation via spark plasma sintering (SPS). The alloy design strategy and detailed fabrication processes are presented. The thermal stability of this DPNC-MEA was investigated and the cor responding mechanism has been proposed.
ball-milling machine. Al, Co, Cu, Ni powders with an equal-atomic ratio (300 mesh, purity > 99.9%, purchased from Aladdin Reagent Co., LTD) and tungsten carbide balls were placed in stainless-steel vials with a powder-to-ball weight ratio of 1:10. Milling was carried out for 50 h in an argon atmosphere with a rotation speed of 300 rpm with a two-step method: firstly, a dry milling for 40 h without a process control agent (PCA) and then wet milling for 10 h with ethanol as a PCA. Subse quently, the mechanical-alloyed AlCoCuNi powder was consolidated by SPS (furnace SE-607, FCT Group, Germany) at 900 � C, with a heating � rate of 100 C/min and a constant pressure of 50 MPa. 2.3. Mechanical and microstructure characterization
2. Material and methods
The DPNC-AlCoCuNi specimens were annealed at a temperature ranging from 500 to 1100 � C for various time to investigate the thermal stability. Then the annealed samples were subjected to Vickers micro hardness test at room temperature with a load of 500 g and a dwell time of 15 s. Each examined sample was indented for at least 8 times in different areas. The compressive tests were carried out using an Instron5967 machine using a strain rate of 1 � 10 3 s 1. The phase composition was determined by X-ray diffraction (XRD, PANalytical X’Pert Pro MRD diffractometer) by using Cu Kα radiations (wavelength λ ¼ 0.154 nm) at a voltage of 45 kV and a tube current of 200 mA. The microstructural characterization was carried out using scanning electron microscope (SEM, SU8020) and transmission electron microscope (TEM, Tecnai G2 F20) with an energy-dispersive X-ray spectrometer (EDS). Thin speci mens for TEM observation were firstly dimpled and then polished using a Gatan 691 precision ion-milling machine with a low energy ion-beam (3 keV).
2.1. The alloy design strategy Since the concept of high-entropy alloys were proposed in 2004 [17], most of the related studies have mainly focused on the design of single-phase high-entropy alloy (HEA) components via maximized configurational entropy to stabilize single-phase structures. Recent re sults show that a high-entropy dual-phase (DP-) HEAs with a high degree of coherent interfaces between two phases exhibit an improved strength as well as excellent ductility, especially at high temperatures, compared to the single-phase HEAs [18]. More importantly, the DP-structure will suppress the recrystallization and retard the grain coarsening even at high temperatures that could improve the thermal stability of materials [19]. Hence, it is promising to explore new DP-HEAs system for the application as high-temperature structural materials. It has been reported that there are several DP-HEA components, such as AlCoCrCu0.5Fe, AlxCoCrFeNi and Al0.5CoCrCuFeNi [20–22]. In these systems with high Fe content, DP-structure is unstable at high temper atures. However, the DP-structure in a Fe-free AlCoCrCuNi alloy can be maintained up to 1200 � C [23]. It is known that the Cr element has a strong tendency to form a brittle σ phase or segregate at grain bound aries during the fabrication process. So in this work, AlCoCuNi medium-entropy alloy (MEA) with Cr element free was designed to avoid the formation of precipitate phases. Besides, to ensure the for mation of stabilized solid-solution phases without additional bulk-metallic glass, empirical conditions of Ω � 1.1 and δ � 6.6% should be satisfied, where Ω is an indicator of the stability of the multicom ponent solid solution and δ is the atomic size difference of the compo nent elements [24], rffiffiffiffiffiffiffi�ffiffiffiffiffiffiffiffiffiffiffiffiffiffi�ffiffiffiffi ri 2 δ ¼ Σci 1 r
3. Results and discussion 3.1. Microstructure and chemical composition of 50 h ball milling powder The SEM image of the as-milled powder and corresponding EDS mapping are shown in Fig. 1. The raw Al, Co, Cu, Ni powder was crushed down and repeatedly welded together to form a uniform distribution of elemental constituent during the first 40 h dry milling. And then a period of 10 h wet milling with ethanol as a PCA is employed to avoid the aggregation of the powders and produce finer particles. After the 50 h milling, powder particles exhibit near-equiaxed and irregular morphology, with particle sizes varying from 5 to 25 μm, as shown in Fig. 1a. Furthermore, the corresponding EDS mapping reveals that each milled powder particle contains homogeneous distribution of Al, Co, Cu and Ni elements as shown in Fig. 1b.
where r ¼ Σci ri is the average atomic radius, ci, ri is the molar content and atomic radius of the ith component, respectively. Tm ¼ Σci ðTm Þi is the average melting point of the alloy system. ΔSmix ¼ RΣci lnci is the is the entropy of mixing for equimolar ratio alloy. ΔH ¼ Σi
3.2. Phase evolution of the AlCoCuNi MEA after MA, SPS, and annealing The manually grinded AlCoCuNi mixture powder with an equalatomic ratio was characterized by X-ray diffraction (XRD) to exhibit Al, Co, Cu, and Ni phases, as shown in Fig. 2a “0 h” sample. After high energy ball milling, diffraction peaks of Cu-based FCC appear and peaks of Co, Cu and Ni are absent, as shown in Fig. 2a “50 h milled” sample. In addition, a new series of diffraction peaks were detected and indexed to the AlCo-based BCC structure, as indicated by red arrows in Fig. 2a. The alloy formation is governed by the inherent thermodynamic and kinetic constraints of the system like binary enthalpy of mixing (ΔHmix) [26]. Phase separation is often observed in HEAs when a high entropy of mixing cannot counteract a significantly positive enthalpy of mixing (or formation of significantly negative enthalpy induced by intermetallic phases). It is suggested that during MA, Co and Ni have a strong ten dency to alloying with Al and produce an AlCoNi-rich BCC phase, owing to the high negative ΔHmix (Fig. 2b), which confirms the XRD results. Also, all peaks of Cu-based FCC and AlCo-based BCC are obviously broadened after ball milling, suggesting a remarkable grain refinement and lattice distortion. These results imply that the formation of
enthalpy of mixing. To construct a thermal stabled dual-phase MEA alloy, the chemical composition of Al25Co25Cu25Ni25 with Ω ¼ 2.22 and δ ¼ 5.82% (using data in Table 1 and Fig. 2b) was designed according to this concept [25]. 2.2. Fabrication methods DPNC-AlCoCuNi MEA powder was synthesized by mechanical alloying (MA) from pure elemental powders in a high-energy planetary
Table 1 Atomic radius, melting temperature of Al, Co, Cu, Ni pure elements. r (nm) Tm (K)
Al
Co
Cu
Ni
0.14317 933
0.12510 1768
0.1278 1358
0.12459 1728
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Fig. 1. (a) SEM image of 50 h as-milled AlCoCuNi powders, and (b) corresponding EDS mapping of Al, Co, Cu, Ni, exhibiting the homogeneous distribution.
Fig. 2. (a) XRD patterns of AlCoCuNi mixture powders before and after high energy ball milling, and SPSed bulk DPNC-AlCoCuNi MEA, (b) the value of enthalpy of mixing (ΔHmix) for atomic pairs [32], (c) the (111) diffraction peak of Cu-rich FCC in bulk DPNC-AlCoCuNi and the relative standard (111) diffraction peak in pure Cu.
Fig. 3. TEM characterizations of (a)–(c) the as SPSed and (d)–(e) 900 � C annealed bulk DPNC-AlCoCuNi MEA: (a, d) bright-field images, (b, e) dark-field images and (f) grain size distributions. 3
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DPNC-AlCoCuNi MEA powder after MA. An effective SPS technique was employed to further consolidate this powder and to get a more equi librium bulk DPNC-AlCoCuNi. The corresponding XRD result of the SPSed bulk sample revealed a more obvious dual-phase structure with Cu-rich FCC (lattice constant, 0.36081 nm) and AlCoNi-rich BCC (lattice constant, 0.28646 nm) phases. The peak position of the FCC phase slightly shifts to the low angle in comparison with pure Cu (Fig. 2c), indicating partial dissolution of larger sized atoms (Al, Table 1) resulting in expansion of FCC lattice. The bright-field and dark-field TEM images of the SPSed AlCoCuNi MEA are presented in Fig. 3a and Fig. 3b, respectively. The SPSed specimens show a NC structure with a mean grain size of about 46 nm. It is worth mentioning that the grain size in both Cu-rich FCC and AlCoNirich BCC was counted as a single species because of their comparable crystalline sizes and distributions. Meanwhile, from the corresponding selected area electron diffraction (SAED) pattern (Fig. 3c), it can be seen that some diffraction rings belong to both FCC and BCC phase which suggest the existence of similar interplanar crystal spaces and forebodes ideal crystallographic orientations between these two phases. After relatively high-temperature annealing at 900 � C for 1 h, the NC structure was maintained and the mean grain size kept at 52 nm, suggesting un precedented thermal stability in our DPNC-AlCoCuNi MEA. The com parable histograms of the grain size distribution (Fig. 3f) in the as SPSed and annealed ones further confirm the high structural stability. The grain size distribution is calculated from more than 300 grains from several TEM images captured from different regions of the samples.
critical temperature of about 900 � C, which is significantly higher than that of other NC materials as indicated by the open blue arrows in Fig. 4b. To further evaluate the thermal stability of DPNC-AlCoCuNi MEA, a series of long-duration annealing up to 50 h was performed, the Vickers hardness dependence of annealing time at various temperatures are shown in Fig. 5. In the case of 900 � C annealing with increasing the annealing time to 10 h, there is just a slight decrease of the harness to 591 HV (ΔHV ~ 8.5%, from the initial 645 HV). It is worth noting that the change in hardness value is negligible for specimens annealed at 900 � C when further increasing the annealing time to 20 h or even to 50 h. For instance, after annealing at 900 � C for 50 h, the hardness of the annealed specimen is still as high as 580 HV. The SEM image of the related fracture surface reveals that the nanostructure of the alloy is still maintained after annealing at 900 � C for 50 h, except for a few grown micron-sized grains as shown in Fig. 5. The newly grown micron-sized grains are Cu-rich FCC structures as illustrated in the following TEM analysis. In the case of 950 � C annealing, a moderate decrease of the harness to 548 HV (ΔHV ~ 15%) is observed after annealing for 10 h. Further increasing the annealing (at 950 � C) time to 50 h, a high level of hardness value of about 508 HV is maintained. The corresponding microstructure of the fracture surface shows that most of the nano structure is retained even after annealing at 950 � C for 50 h, except for some grown micron-sized grains as shown in Fig. 5. While after 10 h annealing at 1000 � C, the hardness is decreased down to 480 HV, about 26% lower than that of the SPSed one, suggesting a typical micro structure instability as indicated in fracture surface morphology shown in the inset of Fig. 5. All these indicate the DPNC-AlCoCuNi MEA indeed possesses an excellent resistance to long-term annealing softening when the annealing temperature is below 900 � C.
3.3. Mechanical properties and stabilities of the DPNC-AlCoCuNi MEA The engineering stress-strain curve of DPNC-AlCoCuNi MEA under uniaxial compression at room temperature is presented in Fig. 4a. The specimen shows an extremely high compressive strength of 2.0 GPa with the strain up to 10.1%, which is higher than that of reported AlCuCr FeMnW0.5(1510 MPa) [27], AlCuNiFeCr (1960 MPa) [28], and Al0.3FeNiCo1.2CrCu (1635 MPa) [29]. It is well known that the hardness is closely related to the microstructure, therefore the structural stability can be further reflected by the evolution of hardness with annealing temperatures. Our DPNC-AlCoCuNi MEA exhibits extremely higher hardness up to 645 HV than that of other NC materials like the NC-CoCrFeMnNi HEA, NC-CoCrFeNi MEA and some Cu-based NC metals and alloys like NC-Cu and Cu/Ta nano-lamellar multilayer, as compared in Fig. 4b. More importantly, the DPNC-AlCoCuNi MEA maintains this high hardness value even after 900 � C annealing for 1 h, suggesting the outstanding thermal stability. Further increasing the annealing tem perature to 1000 � C, the hardness just starts decreasing to 586 HV because of the structure failure of the nanosized Cu-rich FCC. That’s to say, the DPNC-AlCoCuNi MEA could maintain an unprecedented ther mal stabled nano-structure as well as a high strength/hardness until a
Fig. 5. Vickers hardness evolution of DPNC-AlCoCuNi MEA after annealing at 900, 950, and 1000 � C for the time rang of 1–50 h; the SEM images and the inset show the nanostructure of the alloy annealing at 950 � C for 50 h.
Fig. 4. (a) Engineering stress-strain curve of the bulk SPSed DPNC-AlCoCuNi MEA specimen under uniaxial compression at room temperature. (b) Vickers microhardness evolution of DPNC-AlCoCuNi MEA after annealing for 1 h in the temperature range of 500–1100 � C, values of other NC-materials are present for comparison. 4
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Fig. 6 compares the SEM microstructure of the fracture surface after annealing at 900, 950, and 1000 � C for 5 and 10 h. It can be seen that the grain structure consists of nanostructured grains and recrystallized coarse grains (as indicated by the black open arrows in Fig. 6). The proportion of recrystallization under various annealing conditions was quantified using an image recognition algorithm and the results are presented in Fig. 6. A negligible recrystallization proportions of about 3.5% and 4.3% are observed after annealing at 900 � C for 5 h and 10 h, respectively. And the proportion of recrystallization increases with increasing the annealing temperature from 900 � C to 1000 � C, which is consistent with the decreasing trend of hardness. For example, recrys tallization proportion increased to 23.3% after annealing at 1000 � C for 10 h, which is 6 times than of the 900 � C annealed one.
profile confirms this dual-phase nanocrystalline separation structure, as depicted in Fig. 7b. This DPNC separation structure was bred from MA and consolidated during SPS, resulting in low-energy states in both Curich and AlCoNi-rich NCs because of the high entropy effects. In addi tion, these extensively separated PBs could generate a sufficient solutedrag effect, suppress kinetics for coarsening and hence significantly in crease the thermal stabilities. The HRTEM image (Fig. 7c) provides further evidence that Cu-rich and AlCoNi-rich NCs constructed a continuous PB as indicated by the white dotted line and the inset cor responding fast Fourier transform (FFT) images. The representative in verse fast Fourier transform (IFFT) of the yellow dotted square region in Fig. 7c is presented in Fig. 7d to highlight the details of the PB. An evidently special orientation for the FCC and BCC crystals with (31–1)FCC//(01-1)BCC and [011]FCC//[-111]BCC is observed. The present case shows that the orientation is close to the Kurdjumov-Sachs (K–S) relationships [30] but contains an 18� misorientation between the close-packed planes of (1–11)FCC and (110)BCC. In addition, the results calculated from diffraction spots of the PB region exhibit a negligible lattice mismatch (δ ¼ 0.4%) between the parallel planes of (31–1)FCC and (02-2)BCC, suggesting the formation of a coherent low-energy PB. It is interesting to mention that the lattice mismatch between the (31–1)FCC and (02-2)BCC is calculated to be δ~6.9% from the matrix region (away from the PB, as shown in Fig. 7c inset FFT images), which is quite larger than the case of PB region and beyond the coherent scope. However, a slight decrease of lattice space of (31–1)FCC and an increase
3.4. Interface structure of the DPNC-AlCoCuNi MEA To further comprehend essential factors leading to the high thermal stability in this DPNC-AlCoCuNi MEA, extensive interfaces between Curich FCC and AlCoNi-rich BCC phases were detailedly investigated by high-angle annular dark-field scanning transmission electron micro scopy (HAADF-STEM), high-resolution TEM (HRTEM) and EDS. The HAADF-STEM image (Fig. 7a) accentuates the Cu-rich nanograins on the basis of atomic-number contrast, which were separated and surrounded by AlCoNi-rich nanograins in the SPSed DPNC-AlCoCuNi MEA. A locally amplified image with the corresponding HAADF-STEM-EDS line-scan
Fig. 6. SEM microstructures of DPNC-AlCoCuNi MEA after annealing at 900, 950, and 1000 � C for the 5 and 10 h; and the proportion of recrystallization are presented. 5
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Fig. 7. Microstructure of DPNC-AlCoCuNi MEA: (a) a HAADF image, (b) HAADF-STEM-EDS line-scan across three different nanograins, (c) HRTEM image showing the PB, (d) IFFT from the yellow dotted square region in (c). (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
of (02-2)BCC have been triggered through a self-adaptation process to decrease the interface energy and thus a stabled low-energy coherent- state PB has been constructed. This self-adaptation process has been indicated from the atomic redistribution of partial Al atoms (with the largest atomic radius) [31], which diffuse from the Cu-rich FCC to AlCoNi-rich BCC during high-temperature consolidation to reduce the PB interface energy. This Al diffusion could be confirmed by the corre sponding HADDF-STEM-EDS line-scan profile of Al distribution (the blue line in Fig. 7b) which has a gentle slope through the Cu-rich FCC to AlCoNi-rich BCC interface. In addition to these extensive coherent PBs, a few of other lowenergy interfaces like low-angle GBs in Cu-rich FCC were also observed, as presented in Fig. 8. Two Cu-rich FCC grains were viewed along the z ¼ [011]FCC, as indicated by the inset corresponding FFT images. The corresponding HRTEM image (Fig. 8a) shows a low-angle GB in the Cu-rich FCC phase as indicated by the white dotted line. In this case, some lattice distortion, dislocations, and a resulting disorien tation of about 8� were observed, as presented in Fig. 8b. These lowangle GBs with low energy states in Cu-based alloys have been indi cated exhibiting better thermal stability than monolithic nanocrystals with high-angle grain boundaries [29]. Combining the XRD analysis, the chemical composition, and the microstructure characterization, the formation process of this thermally stabled DPNC-AlCoCuNi MEA was summarized, as illustrated schemat ically in Fig. 9. After 50 h ball milling, Al, Co, Cu and Ni pure elements with an equal-atomic ratio roughly constructed the DP structure (phase separation) with both the AlCoNi-rich BCC and Cu-rich FCC phases, as indicated by the blue lines in Fig. 9b. In addition, remarkable grain re finements to form nanocrystalline grains/cells (indicated by the black dotted lines in Fig. 9b) could be confirmed by the obviously broadened XRD diffraction peaks as shown in Fig. 2a. Finally, these extensive nanosized PBs between the Cu-rich FCC to AlCoNi-rich BCC phases and some nano-sized low-angle GBs were preserved and further consolidated
(reducing the lattice defects like vacancies and dislocations) to form the stabled bulk DPNC-AlCoCuNi MEA during the effective SPS process, as shown in Fig. 9c. Hence, it is comprehensible that the bulk DPNCAlCoCuNi MEA achieves the unprecedented thermal stability: i) MA constructing the DPNC MEA powder with Cu-rich FCC and AlCoNi-rich BCC NCs with both the high entropy and sluggish diffusion effects, ii) high-temperature SPS consolidating DPNC structure through a selfadaptation process, resulting in extensive thermal-stabled low-energy coherent PBs and some low-angle GBs. As a representative example, the microstructure after 1000 � C annealing for 1 h is shown in Fig. 10. The DPNC-AlCoCuNi MEA is hard to maintain the nanostructure due to the instability of Cu-rich FCC. Correspondingly, the Cu-rich FCC nanocrystalline grains begin to grow into micron ones as highlighted by the white dotted line in Fig. 10. It is worth mentioning that AlCoNi-rich BCC nanocrystalline grains are still retained even after the 1000 � C annealing. 4. Conclusions A bulk DPNC-AlCoCuNi MEA with unprecedented thermal stability was successfully processed via MA and SPS consolidation. This NC MEA has an average grain size of 46 nm, an extremely high compressive strength of 2.0 GPa with the strain up to 10.1% and a high Vickers microhardness of 645 HV. More importantly, this DPNC-AlCoCuNi material could maintain the nanostructures as well as a high hardness of about 580 HV even after annealing at 900 � C for 50 h. This critical temperature of about 900 � C is significantly higher than that of other Cu containing NC materials and NC MEA alloys [15, 16, 27, 31–34]. The extremely high thermal stability has been indicated from i) the thermalstabled matrix including Cu-rich FCC and AlCoNi-rich BCC NCs with high entropy and sluggish diffusion effects and ii) extensive thermalstabled low-energy coherent PBs and some low-angle GBs.
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Fig. 8. (a) High-resolution TEM (HR-TEM) of the grain boundary, (b) Inverse fast Fourier transform (IFFT) image from the region indicated by the yellow dotted square in (a). The fast Fourier transform (FFT) images in the inset of (a) show the two FCC grains. The white dotted line indicates the grain boundary. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
Fig. 9. A Schematic diagram showing the stable DPNC-AlCoCuNi MEA formation process.
Acknowledgements This work was financially supported by the National Natural Science Foundation of China (Grant Nos.: 51771184, 11735015, 11575241, 51801203, 11674319), the Natural Science Foundation of Anhui Prov ince (Grant No. 1808085QE132) and the Open Project of State Key Laboratory of Environment friendly Energy Materials (Grant No. 18KFHG02). References [1] C.A. Schuh, T.G. Nieh, H. Iwasaki, The effect of solid solution W additions on the mechanical properties of nanocrystalline Ni, Acta Mater. 51 (2) (2003) 431–443. [2] K.S. Kumar, H. Van Swygenhoven, S. Suresh, Mechanical behavior of nanocrystalline metals and alloys11The golden jubilee issue—selected topics in materials science and engineering: past, present and future, in: S. Suresh (Ed.), Acta Mater. 51 (19) (2003) 5743–5774. [3] Y.F. Shen, L. Lu, Q.H. Lu, Z.H. Jin, K. Lu, Tensile properties of copper with nanoscale twins, Scr. Mater. 52 (10) (2005) 989–994. [4] K. Lu, Stabilizing nanostructures in metals using grain and twin boundary architectures, Nat. Rev. Mater. 1 (5) (2016). [5] L.F. Zeng, P. Fan, L.F. Zhang, R. Gao, Z.M. Xie, Q.F. Fang, X.P. Wang, D.Q. Yuan, T. Zhang, C.S. Liu, He irradiation effects in bulk Cu/V nanolayered composites fabricated by cross accumulative roll bonding, J. Nucl. Mater. 508 (2018) 354–360. [6] H.J. Fecht, Intrinsic instability and entropy stabilization of grain boundaries, Phys. Rev. Lett. 65 (5) (1990) 610. [7] B. Günther, A. Kumpmann, H.D. Kunze, Secondary recrystallization effects in nanostructured elemental metals, Scr. Metall. Mater. 27 (7) (1992) 833–838. [8] C.C. Koch, R.O. Scattergood, K.A. Darling, J.E. Semones, Stabilization of nanocrystalline grain sizes by solute additions, J. Mater. Sci. 43 (23–24) (2008) 7264–7272. [9] Z. Lei, X. Liu, H. Wang, Y. Wu, S. Jiang, Z. Lu, Development of advanced materials via entropy engineering, Scr. Mater. 165 (2019) 164–169. [10] Y. Zou, H. Ma, R. Spolenak, Ultrastrong ductile and stable high-entropy alloys at small scales, Nat. Commun. 6 (2015) 7748. [11] K.Y. Tsai, M.H. Tsai, J.W. Yeh, Sluggish diffusion in Co–Cr–Fe–Mn–Ni high-entropy alloys, Acta Mater. 61 (13) (2013) 4887–4897. [12] N. Zhou, T. Hu, J. Huang, J. Luo, Stabilization of nanocrystalline alloys at high temperatures via utilizing high-entropy grain boundary complexions, Scr. Mater. 124 (2016) 160–163. [13] O. Anderoglu, A. Misra, H. Wang, X. Zhang, Thermal stability of sputtered Cu films with nanoscale growth twins, J. Appl. Phys. 103 (9) (2008), 094322.
Fig. 10. A TEM image showing the 1000 � C annealed DPNC-AlCoCuNi MEA.
Declaration of competing interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. CRediT authorship contribution statement H.W. Deng: Conceptualization, Methodology, Investigation, Data curation, Writing - original draft. Z.M. Xie: Methodology, Formal analysis, Data curation, Writing - review & editing. M.M. Wang: Investigation. Y. Chen: Validation. R. Liu: Resources, Supervision. J.F. Yang: Formal analysis, Data curation. T. Zhang: Project administration, Funding acquisition, Writing - review & editing. X.P. Wang: Supervi sion. Q.F. Fang: Project administration, Funding acquisition. C.S. Liu: Project administration, Funding acquisition. Y. Xiong: Funding acquisition.
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Materials Science & Engineering A 774 (2020) 138925 [24] Y. Zhang, T.T. Zuo, Z. Tang, M.C. Gao, K.A. Dahmen, P.K. Liaw, Z.P. Lu, Microstructures and properties of high-entropy alloys, Prog. Mater. Sci. 61 (2014) 1–93. [25] Z. Fu, L. Jiang, J.L. Wardini, B.E. MacDonald, H. Wen, W. Xiong, D. Zhang, Y. Zhou, T.J. Rupert, W. Chen, E.J. Lavernia, A high-entropy alloy with hierarchical nanoprecipitates and ultrahigh strength, Sci. Adv. 4 (10) (2018), eaat8712. [26] M. Vaidya, A. Prasad, A. Parakh, B.S. Murty, Influence of sequence of elemental addition on phase evolution in nanocrystalline AlCoCrFeNi: novel approach to alloy synthesis using mechanical alloying, Mater. Des. 126 (2017) 37–46. [27] D. Kumar, V.K. Sharma, Y.V.S.S. Prasad, V. Kumar, Materials-structure-property correlation study of spark plasma sintered AlCuCrFeMnWx (x ¼ 0, 0.05, 0.1, 0.5) high-entropy alloys, J. Mater. Res. 34 (5) (2019) 767–776. [28] A.I. Yurkova, V.V. Cherniavsky, V. Bolbut, M. Krüger, I. Bogomol, Structure formation and mechanical properties of the high-entropy AlCuNiFeCr alloy prepared by mechanical alloying and spark plasma sintering, J. Alloy. Comp. 786 (2019) 139–148. [29] S. Yang, J. Pi, W. Yang, H. Zhou, D. Zhu, Deformation twinning structure and interface in a FCC-based Al0.3FeNiCo1.2CrCu high-entropy alloy matrix composites, Mater. Lett. 214 (2018) 50–52. [30] J.H. van der Merwe, G.J. Shiflet, The role of structural ledges at phase boundaries—III. F.C.C.-B.C.C. interfaces in Kurdjumov-Sachs orientation, Acta Metall. Mater. 42 (4) (1994) 1199–1205. [31] L.J. Santodonato, Y. Zhang, M. Feygenson, C.M. Parish, M.C. Gao, R.J.K. Weber, J. C. Neuefeind, Z. Tang, P.K. Liaw, Deviation from high-entropy configurations in the atomic distributions of a multi-principal-element alloy, Nat. Commun. 6 (2015) 5964. [32] A. Takeuchi, A. Inoue, Classification of bulk metallic glasses by atomic size difference, heat of mixing and period of constituent elements and its application to characterization of the main alloying element, Mater. Trans. 46 (12) (2005) 2817–2829.
[14] Y. Chen, Y. Liu, F. Khatkhatay, C. Sun, H. Wang, X. Zhang, Significant enhancement in the thermal stability of nanocrystalline metals via immiscible triphases, Scr. Mater. 67 (2) (2012) 177–180. [15] E. Botcharova, J. Freudenberger, L. Schultz, Mechanical and electrical properties of mechanically alloyed nanocrystalline Cu–Nb alloys, Acta Mater. 54 (12) (2006) 3333–3341. [16] L.F. Zeng, R. Gao, Q.F. Fang, X.P. Wang, Z.M. Xie, S. Miao, T. Hao, T. Zhang, High strength and thermal stability of bulk Cu/Ta nano lamellar multilayers fabricated by cross accumulative roll bonding, Acta Mater. 110 (2016) 341–351. [17] J.W. Yeh, S.K. Chen, S.J. Lin, J.Y. Gan, T.S. Chin, T.T. Shun, C.H. Tsau, S.Y. Chang, Nanostructured high-entropy alloys with multiple principal elements: novel alloy design concepts and outcomes, Adv. Eng. Mater. 6 (5) (2004) 299–303. [18] K.R. Lim, K.S. Lee, J.S. Lee, J.Y. Kim, H.J. Chang, Y.S. Na, Dual-phase high-entropy alloys for high-temperature structural applications, J. Alloy. Comp. 728 (2017) 1235–1238. [19] J. Hou, X. Shi, J. Qiao, Y. Zhang, P.K. Liaw, Y. Wu, Ultrafine-grained dual phase Al0.45CoCrFeNi high-entropy alloys, Mater. Des. 180 (2019) 107910. [20] W.-R. Wang, W.-L. Wang, S.-C. Wang, Y.-C. Tsai, C.-H. Lai, J.-W. Yeh, Effects of Al addition on the microstructure and mechanical property of AlxCoCrFeNi highentropy alloys, Intermetallics 26 (2012) 44–51. [21] C.W. Tsai, M.H. Tsai, J.W. Yeh, C.C. Yang, Effect of temperature on mechanical properties of Al0.5CoCrCuFeNi wrought alloy, J. Alloy. Comp. 490 (1) (2010) 160–165. [22] K.E. Guan-Yu, Swe-Kai CHEN, Tung HSU, Jien-Wei YEH, FCC and BCC equivalents in as-cast solid solutions of Al x Co y Cr z Cu 0.5 Fe v Ni w high-entropy alloys, Proc. Prehist. Soc. 31 (2) (2006), 633-381. [23] Ł. Rogal, On the microstructure and mechanical properties of the AlCoCrCuNi high entropy alloy processed in the semi-solid state, Mater. Sci. Eng. A 709 (2018) 139–151.
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