Accepted Manuscript Ageing behaviour of equiatomic Al20Co20Cu20Ni20Zn20 high entropy alloy
consolidated
S. Mohanty, N.P. Gurao, P. Padaikathan, Krishanu Biswas PII: DOI: Reference:
S1044-5803(17)30636-8 doi: 10.1016/j.matchar.2017.04.011 MTL 8634
To appear in:
Materials Characterization
Received date: Revised date: Accepted date:
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ACCEPTED MANUSCRIPT Ageing Behaviour of Equiatomic Consolidated Al20Co20Cu20Ni20 Zn20 High Entropy Alloy S. Mohanty1, N. P. Gurao1, P.Padaikathan2 and Krishanu Biswas1#
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Department of Materials Science and Engineering, Indian Institute of Technology Kanpur, Kanpur -208016, INDIA Department of Materials Engineering, Indian Institute of Science, Bangalore -560012, INDIA
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Abstract
The effect of ageing treatment on the evolution of microstructure and hardness of
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consolidated Al20Co20Cu20Ni20Zn20 high entropy alloy is reported. The consolidated alloy specimens were solutionized by heat treatment at 1433 K for 96 h. The solutionized specimens were subsequently subjected to ageing treatment at 673, 773, 873 and 1073 K for
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48 h to investigate the effect of ageing treatment on the evolution of microstructure as well as micro-hardness of the alloy. The fully solutionized specimens reveal formation of single phase FCC phase. The artificially aged alloy consists of Ni and Al rich FCC α (L12)
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precipitate in FCC (γ) matrix. This is akin to precipitation in Ni-based superalloys. The
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hardness measurement of the aged specimens indicate peak age-hardening (~6.2 GPa) at 773 K. Therefore, the present study demonstrates possibility of age hardening of consolidated
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alloys akin to sinter ageing.
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Keywords: HEA, Ageing, Precipitation hardening, Spark plasma sintering, Hardness, Multicomponent alloys
#Corresponding author: email:
[email protected], Phone: +91-512-2596184, Fax: +91-512-2597550
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ACCEPTED MANUSCRIPT 1. Introduction The design strategy of almost all traditional alloy systems is usually based on one major constituent, such as aluminium, iron, copper, titanium and magnesium based alloys along with the addition of other minor elements in order to enhance the microstructure, properties and performances of the alloy systems [1, 2]. According to the traditional concept, a multicomponent alloy system would lead to the formation of many intermetallic compounds or
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complex structures, which not only increase the complications in analyzing the microstructure but also have an adverse effect on the mechanical properties of the alloy
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system. To overcome this constraint, Yeh et al. [3] have proposed the concept of “high entropy alloys” (HEAs) for multi-component alloy design consisting of at least five principal
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metallic elements with the concentration of each elements in equi-molar or near equi-molar ratios lying between 5 at.% and 35 at.%. It has attracted extensive attention worldwide by
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various researchers [3]. The high entropy of mixing (≥ 1.5R) arising in the form of different configurations of arranging atoms leads to formation of simple solid solution structures such
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as FCC (face-centered cubic), BCC (body-centred cubic) and HCP (hexagonal close-packed) instead of intermetallics and complex microstructures in these multi-component multiprincipal HEAs [4-7] contributing to sluggish atomic diffusion [8], large lattice distortion [9]
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and cocktail effect [10]. These core effects play important role in enhancing the formation
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and stabilization of the phases with simple crystal structure along with easy formation of nano precipitates [4, 11]. As a result, HEAs exhibit many attractive properties including high hardness and plausible resistance to anneal softening, good wear, oxidation and corrosion
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resistance [4, 12-15].
Recent research activities regarding HEAs mainly concentrate on the design and development It not only depends on the variety of principal elements and their
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of novel HEAs.
composition, but also on the various processing routes such as arc melting, melt spinning technique, mechanical alloying followed by sintering because phase formation and microstructural evolution strongly depends on processing route for a particular HEA [16-21]. Thus, it opens up large number of possibilities of microstructural design and development of materials with novel properties. One such novelty involves precipitation hardening in the high entropy alloys [22-28]. The present investigation demonstrates that it is possible to adopt artificial ageing treatment for equiatomic Al20Co20Cu20Ni20Zn20 alloy to form L12 ordered precipitates in a FCC matrix, similar to Ni-based superalloys. Therefore, it provides new vistas of alloy design and thus opens up novel processing of HEAs. 2
ACCEPTED MANUSCRIPT It is to be noted that the previous study has reported the sinter ageing of Al20Co20Cu20Ni20Zn20 HEA [18], synthesized by high energy ball milling followed by consolidation by advanced spark plasma sintering (SPS) at different sintering temperature at a heating rate of 100 K/min and applied pressure of 50 MPa. The alloy showed single phase supersaturated FCC solid solution in the mechanically alloyed (MA) powders. However, the consolidation of the MA powder using SPS led to the precipitation of near cuboidal shaped
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ordered FCC α (L12) phase within the grains of FCC γ. The room temperature hardness measurement of the sintered alloys suggests age hardening of the as-milled powder during
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sintering. The hardness of the sintered alloys increases with increase in temperature and has the highest value 6.3 GPa at 1073 K followed by a decrease in hardness at higher sintering
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temperatures.
To expand the scope of understanding regarding microstructural evolution as well as the age-
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hardening phenomena of the sinter-aged Al20Co20Cu20Ni20Zn20 HEA, the present investigation deals with the artificial ageing behaviour of the consolidated sinter-aged
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Al20Co20Cu20Ni20Zn20 HEA, by solutionizing the sample at 1433 K for 96 h. This was followed by ageing treatment at various temperatures (673, 773, 873, 1073 and 1273 K) for 48 h respectively, in order to control the size, shape and distribution of the precipitates in the
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alloy. In addition to the aforementioned isochronal ageing treatment, isothermal ageing was
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also carried out at 773 K for different time period. The age hardening behaviour of the consolidated Al20Co20Cu20Ni20Zn20 HEA is investigated in order to optimize the mechanical properties. The effect of the ageing condition on the mechanical properties of the solutionized
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HEA is evaluated using hardness measurements. Finally, the mechanical properties of the alloy are correlated with the microstructural evolution in the light of available literature.
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2. Experimental procedures Al20Co20Cu20Ni20Zn20 HEA was synthesized by ball milling high purity commercial Al powder (mean particle size ~ 44 μm, 99.8% purity, Alfa-Aesar), Co powder (mean particle size < 150 μm, 99.9% purity, Sigma-Aldrich), Cu powder (mean particle size ~ 3 μm, 99.7% purity, Sigma-Aldrich), Ni powder (mean particle size < 150 μm, 99.99% purity, SigmaAldrich), Zn powder (≥ 99% purity, Sigma-Aldrich) for 15 h in dry condition, at a speed of 200 rpm with ball to powder weight ratio (BPR) of 20:1 under high purity argon atmosphere. Milling of the elemental powders was carried out in Pulverisette-5 planetary ball mill (Fritsch, Germany) with tungsten carbide (WC) vials and balls. Prior to the milling, the 3
ACCEPTED MANUSCRIPT hermetically sealed WC vials containing powder and balls were evacuated to 10-3 torr and refilled with high purity (99.9%) argon gas in order to avoid oxidation of the powders. After every 1 hour of milling, the milling operation was intermittently stopped and after properly evacuating the vials, high purity argon gas was subsequently replenished into the sealed vials. The spark plasma sintering (SPS) of powders was carried out at 1223 K using a Dr. Sinter 515S apparatus (SPS Syntex INC., Kanagawa Japan) with a pulse on-off ratio of 12:2. The as milled powder loaded in a 15 mm diameter graphite die was placed inside the SPS chamber
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between two graphite electrodes under a vacuum level of 6 103 torr. High purity (99.9%)
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argon gas was then purged in the chamber with a flow rate of 2 litre/min. to ensure minimal oxidation of the powder. The powder was heated with the heating rate of 100 K/min. to 1223
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K and held with a uniaxial pressure of 50 MPa throughout the sintering cycle at the sintering temperature for 6 min before it was furnace cooled to room temperature. The temperature of
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the sample was monitored using a K1-type chromel-alumel thermocouple, which was inserted in the die at a distance of 0.002 m from the inner die so that the recorded temperature
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difference between the sample and the thermocouple tip was minimum. During the SPS experiment, while a constant voltage of 20V was applied, the current flow used was around 900 Amp. Then solution annealing of the sintered specimens were performed at 1433 K for
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96 h by sealing the sintered specimens in evacuated quartz tubes and quenching them in ice
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cold water after 96 h. The quartz tube were thrice evacuated to 10-6 mbar and simultaneously purged with argon gas and finally filled with argon gas before sealing. The solutionized specimens were again sealed inside different quartz tubes as mentioned above and then were
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subjected to ageing treatment at different temperatures (673, 773, 873, 1073, and 1273 K) for 48 h followed by ice water quenching at room temperature. Few solutionized samples were sealed and subjected to isothermal ageing treatment at 773 K for different time period ranging
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from 2-96 hour.
The phase identification of the polished and flat bulk as well as the ball-milled powder samples was performed using X-ray diffractometer (XRD, Bruker D8 FOCUS diffractometer, Germany) in the θ-2θ configuration with V-filtered Cr Kα (λ = 0.228976 nm) radiation. The voltage and current applied to the Cr target were 40 kV and 40 mA, respectively. The sample was rotated at a speed of 15 rpm to prevent the effect of preferred orientation. Fine scale microstructural observations were carried out using a 200 KV transmission electron microscope (TEM, FEI Tecnai G2 UT-20, The Netherland). Samples for TEM observation were prepared by the standard preparation technique, which includes cutting of 3-mm disks 4
ACCEPTED MANUSCRIPT of sintered samples, grinding up to 60-70 μm thick by mechanical polishing using emery paper. As a next step, the thickness of the polished sample was further reduced to 20 μm by dimpling using Gatan, model 656, precision dimpler grinder, USA rotating at a speed of 4 rpm. Thereafter, the sample was subjected to a precision ion polishing system (PIPS: model: 691, Gatan, USA), using argon gas ions at 5 keV and initially at higher angle of 5° until perforation occurred and later at a lower angle of 2° to obtain the electron transparent region.
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The differential scanning calorimetric (DSC) measurements on the samples were carried out using DSC Model STA 449 Fe Jupiter, Germany. Sample was heated and cooled at a rate of 10°C using alumina pan with base line correction incorporating correction mode. Micro-
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hardness of the samples were measured using Vickers micro hardness tester (Digi-test, VTP-
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6046, Bareiss Prüfgerätebau, GmbH, Germany) with a load of 100 g and dwell time of 10s. In order to obtain a statistically pertinent data, at least fifty measurements were carried out for
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each sample and then the average value was reported with error bars, indicating standard deviation. The homogeneity of the SPS pellet was confirmed by taking hardness
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measurement on both sides of the SPS pellet for all the heat treatment conditions.
3. Results
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3.1. Differential Scanning Calorimetric (DSC) Measurement
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A detailed thermal characterization of the consolidated samples has been carried to obtain the important temperatures for solutionizing as well as ageing treatment. The typical DSC thermograms (heating and cooling) for the equiatomic consolidated Al20Co20Cu20Ni20Zn20
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HEA are shown in Figure 1. It is evident from the heating curve that the precipitates formed in the consolidated sample during sintering starts dissolving at 973 K and finished at 1443 K.
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Thereafter the sample starts melting at about 1498 K. Therefore, the consolidated samples were solutionized at 1433 K. The thermogram during cooling reveals presence of four distinct peaks at 1198, 1153, 1053 and 973K. These peaks correspond to the precipitates forming in the fully solutionized specimen during cooling. It is to be mentioned here that precipitation generally occurs through multiple steps comprising of cluster formation followed by formation of coherent, semi-coherent and finally incoherent precipitates [19]. It is expected that the four peaks in the DSC thermogram may correspond to the formation of different stages of a single precipitate obtained in the present alloy as discussed posteriori. In the following, the results obtained during thermal analysis have been utilized for ageing treatment of the equiatomic HEA. 5
ACCEPTED MANUSCRIPT 3.2. Phase analysis Figure 2 shows the X-ray diffraction (XRD) patterns of Al20Co20Cu20Ni20Zn20 HEA sintered at 1223 K, solutionized sample and aged alloys at different ageing temperatures. The sintered specimen reveals the formation of (FCC) and α (L12) phases. The XRD pattern of sintered specimen after solution treatment at 1433 K for 96 h is characterized by the absence of diffraction peaks corresponding to α (L12) phase. This fully solutionized sample has been
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used for further ageing treatment at different temperatures (673, 773, 873, 1073 and 1273 K) for 48 h. XRD patterns in Figure 2 reveals that no ageing occurred at 1073 K, whereas the
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ageing treatment at 673, 773 and 873 K clearly reveals the diffraction peaks of α (L1 2) phase along with γ (FCC) phase, indicating the possible formation of the precipitates in these
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samples. It is to be noted that and phases are similar with slight variation in the lattice
3.4. Microstructural observations
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parameter.
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We have carried out detailed microstructural characterization of the solutionized and aged specimens. Some representative micrographs will be reported here to illustrate the correlation
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between the processing conditions with microstructure. First, we shall describe the microstructure of the solutionized specimen. This will be followed by the microstructural
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investigation of the aged specimens. Figure 3a shows the back scattered electron (BSE) SEM micrographs of the solutioned specimen, revealing single phase microstructure, devoid of any precipitate in grains. It is to be noted here that BSE SEM micrographs illustrate
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compositional contrast and the phases can distinctly be observed to the best of the resolution of SEM. BSE SEM micrographs for the aged specimens are shown in Figures 3(b) and 3(c).
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The specimen aged at 873 K reveals (Figure 3(b)) grains with speckle like contrast indicating the presence precipitates. Figure 3(c) shows micrographs for specimen aged at 1073 K. The two-phase microstructure is distinctly observed. Thus, the formation of precipitate within grains are observed during ageing treatment at higher temperature. We shall now describe TEM observations. Figure 4(a) shows typical low magnification bright field micrograph of the specimen solutionized at 1433 K for 96 h, revealing the micron sized grains of the β phase. Figure 4(b) shows a high magnification dark field micrograph of one such grain devoid of any precipitate. The selected area diffraction (SAD) pattern obtained from the micron sized grains is shown in Figure 4 (c). The SAD pattern can consistently be 6
ACCEPTED MANUSCRIPT indexed due to the β phase. No superlattice reflections corresponding to α phase can be detected. It is to be noted that solutionization treatment led to extensive grain growth. Figure 5 shows the typical bright field micrographs of the specimen sintered at 773 , 873 and 1073K, revealing formation of precipitate of α within the γ grains. The average precipitate measured at different sintering temperature is shown in Table I. The precipitates grow extensively with temperature. At 773 K, the average size is 12±1 nm and size becomes
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224±25 nm at 1073 K. The precipitates show faceted morphology. One can observe the presence of Moiré fringes in the precipitates, indicating possible lattice coherency between
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the precipitate and matrix. The precipitates are marked by arrows. Figure 5(a) reveals the
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presence of fine scale (5-8 nm) precipitates in the matrix. The number density of precipitate is also found to be small at 773K. Ageing treatment at higher temperature leads to increase of number density as well as the size of the precipitates. Figure 5 (c) shows TEM observation
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of the specimen aged at 1073 K. It indicates the presence of large number of faceted α precipitates within γ grains. It can be observed that a well grown α precipitates with size
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varying from 146 nm to 270 nm within the γ phase matrix. The size of the precipitates has been measured using bright TEM micrograph and the results are shown in Table I, which reveals the precipitate size increases as the sintering temperature is increased. Figure 5(d)
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shows typical selected area diffraction (SAD) pattern obtained from a grain. The composite
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SAD pattern obtained from both γ and α phase is shown in Figure 5 (e). The fundamental diffraction spots confirm that the matrix is FCC (γ) phase, whilst additional weak spots observed in the SAD pattern affirm the presence of α precipitates, having superlattice L12
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structure. It is to be noted that the precipitate size increases as the ageing temperature increases from 673 to 1073 K. The matrix grains show the strain contrast (Figure 5 (e)),
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revealing the presence of defects which may have formed during ageing. The detailed compositional analysis performed using energy dispersive spectroscopic (EDS) analysis of the precipitate and matrix has been performed. Table II shows the results. The precipitates are enriched in Al, Ni and Co whereas the matrix is lean in Al but rich in Zn, Ni, Co, Cu Therefore, it is evident that formation of nano-sized L12 (Ni, Co)3Al type α coherent precipitates within the disordered FCC (γ) matrix plays an important role in age-hardening. This is a typical characteristic of age hardening in the metallic systems [29].
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ACCEPTED MANUSCRIPT 3.5 Age-hardening curve Figure 6(a) shows the age-hardening curve of the alloy obtained after ageing treatment at different temperatures. The hardness of the solutionized specimen (~1.67 ± 0.39 GPa) increases steadily up to 673 K followed by a rapid increase at 773 K. The optimal age hardening occurs at 773 K reaching peak hardness of ~ 6.24 ± 0.47 GPa. Subsequently, there is decrease in hardness for sample obtained at higher ageing temperatures because of
effect of ageing time for the peak aged sample.
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overageing. Isothermal ageing experiments were also carried out at 773 K to understand the Figure 6(b) shows the results of the
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isothermal ageing experiment. It is observed that the hardness increases with ageing time and
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peak hardness is observed for a time of 48 hour following which there is a decrease in the hardness. The age hardening phenomena of sinter-aged Al20Co20Cu20Ni20Zn20 HEA is
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discussed in the light of microstructural and textural evolution in the next section.
4. Discussion Al-containing
multi-component
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The present investigation clearly demonstrates the ageing behaviour of the novel sinter-aged equimolar
Al-Co-Cu-Zn-Ni
alloy.
The
underlying
mechanisms contributing to conventional ageing as well as the previously reported sinter
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ageing [18] for the alloy are discussed below. The differential scanning calorimetry (DSC) curve showing a broad endothermic peak appearing at around 1443 K corresponding to the dissolution of α phase present at the grain
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boundary and within the γ grains respectively. The consolidated Al20Co20Cu20Zn20Ni20 HEA is solution treated at 1433 K for 96 h as this temperature can be considered above the solvus temperature for the alloy. The solutionized sample was then rapidly quenched with the
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resulting microstructure revealing only supersaturated solid solution β (Figure 2 and 3 (a, b)). The absence of α phase in solution treated specimens is attributed to insufficient atomic diffusion leading to complete dissolution of into [30]. In addition, at the solutionizing temperature, the effect of entropy of mixing is large enough to dissolve the α phase in the FCC solid solution. When the solutionized specimen are once again heat treated to an ageing temperature below the solvus temperature, the diffusional processes
are enhanced [30]
causing the precipitation of α phase within γ grains. The crystal structure of the aged alloy represented by the XRD patterns shown in Figure 2 reveal a mixture of a major FCC (γ) solid solution matrix and a secondary α phase with L12 structure after ageing at 873 K. The α 8
ACCEPTED MANUSCRIPT phase appears as a distribution of nearly cuboidal shape fine particles within the γ phase grains in the specimen aged at 773 K, as shown in Figure 5 (b). At lower temperature the effect of mixing entropy is reduced leading to reduced solution limit and the precipitation of stable α phase. An increase in the ageing temperature to 873 K and more intensifies the diffusion process and as a result, α precipitates show coarsening along with increase in aspect ratio of cuboidal shape precipitates. The hardness response of the Al20Co20Cu20Zn20Ni20 HEA
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to isochronal ageing at various temperatures is shown in Figure 6. A substantial increase in the hardness was observed after ageing at 773 K and the hardening essentially disappears at
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1073 K.
Traditionally, hardening mechanisms in polycrystalline FCC materials are as a result of :
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solid solution hardening, grain boundary hardening, dislocation hardening and precipitation
H H 0 H ss H gb H dis H ppt
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hardening [31]. The hardness of the material can be expressed as [32-34]: (1)
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where H0 is the intrinsic hardness, or the so-called lattice friction hardness, and ∆Hss, ∆Hgb, ∆Hdis and ∆Hppt are hardness contributions from solid solution, grain boundary, dislocations, and precipitates, respectively. The hardness resulting from solid solution strengthening by
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using standard model for substitutional solid solution strengthening based on dislocation-
3
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solute elastic interactions is given by [33, 34]: G *s 2 * c H ss 3 * 700 2
1
2
(2)
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3
where G is the shear modulus of the alloy, c is the mole fraction of solute and 33/2is an
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approximate conversion factor from shear stress to hardness. The interaction parameter sis defined as [33, 34]:
s
G
1 0.5 G
3 *
(3)
which combines the effects of elastic and atomic size mismatch, i.e. G and a respectively. The mismatch parameters are defined as:
G
1 G G c
(4)
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ACCEPTED MANUSCRIPT a
1 a a c
(5)
where a is the lattice constant of the alloy. The grain boundary strengthening can be defined using the Hall-Petch relationship [33, 35, 36]: k d
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H gb C
(6)
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where C is a constant, k is the Hall-Petch constant and d is the grain size.
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The dislocation hardening using Bailey-Hirsch relationship is described as [31, 32, 36]: (7)
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H dis CMGb
where C is a constant, M is the Taylor factor that converts shear stress to normal stress for a
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FCC polycrystalline matrix, α is a factor depending on the character of dislocations, G is the shear modulus, b is the Burgers vector and ρ is the dislocation density.
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The precipitation hardening (∆Hppt) is governed by, either through Orowan dislocation bypass mechanism, when the precipitate is incoherent with the matrix or the radius of the
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precipitates exceed a critical value in spite of coherent or dislocation shearing mechanism, in which the precipitates are coherent and their sizes are sufficiently small. [31, 37, 38].
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For the Orowan dislocation bypassing mechanism, the hardness increment is approximated as
0.81mb
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[36]: H ppt CM
2 (1 )
1
ln( 2
ds 2 )(0.615d s d s ) 1 b 3 fs
(8)
where μm is the shear modulus of the matrix, ν the Poisson’s ratio, ds the diameter of the cross-section of the precipitates and fs the volume fraction of the precipitates. It is to be mentioned here that strengthening contribution from individual hardening mechanismns enumerated in equations 2- 8 are for alloys with two or three component system and are expected to be modified considerably in the multi-component HEAs.
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ACCEPTED MANUSCRIPT It is expected that the first four terms namely strengthening from lattice friction stress, solid solution formation, grain size and dislocation density contribute to high hardness of the MA and solutionized sample. HEAs are multi-component multiprincipal elements in equimolar or near-equimolar ratio. The solid solution hardening is mainly due to the atomic size difference from different kind of atoms, which increases the lattice distortion and implies a larger resistance to slip deformation and thus increasing the hardness of the alloy [39, 40].
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Traditionally BCC metals and alloys have significant contribution to overall strength from lattice friction stress while, there is negligible contribution of the friction stress in FCC materials [38]. However, the severe lattice distortion in HEAs with FCC structure leads to
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contribution to overall strength from the lattice friction. In the Al20Co20Cu20 Ni20Zn20 HEA,
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with Al (1.43 Å) and Zn (1.4 Å) have the largest atomic radius in comparison with Co (1.25Å), Cu (1.28 Å) and Ni (1.25 Å) and the differences in atomic size contributes to
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significant lattice distortion and hence friction stress. Similarly, the solid solution hardening in Al20Co20Cu20 Ni20Zn20 HEA can be principally attributed to Al and Zn. In addition, to solid solution hardening effect in MA Al20Co20Cu20 Ni20Zn20 HEA forming supersaturated single
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phase FCC (β) solid solution, grain boundary strengthening also contribute to the hardness of the alloy. In addition, presence of large amount of defects, e.g. vacancies, grain boundaries,
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stacking faults, twins, dislocations, etc., occurred during the MA process of the alloy due to severe deformation conditions [41] also contribute to hardening. Grain-size refinement
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obtained during MA increases the volume fraction of grain boundaries which impede the motion of dislocation and increases the hardness of the alloy [31]. Also the severe plastic deformation during MA forms dislocations and these mobile dislocations interact with each
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other and impede their own motion resulting in dislocation hardening [31, 41]. Solid solution hardening, grain boundary hardening and dislocation hardening are also considered as
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hardening mechanism in the consolidated Al20Co20Cu20 Ni20Zn20 HEA. But, the hardness of MA HEA (~2 ± 0.19 GPa) exhibits a threefold increase in hardness in the consolidated specimens (~6.3 ± 0.03 GPa) [18]. This sharp increase in hardness of the consolidated specimens is attributed to the precipitation hardening caused by the precipitation of ordered FCC (Co, Ni)3Al rich α phase within disordered FCC (γ) matrix, which are absent in MA HEA as observed by Mohanty et al. [18]. Thus, the high hardness of the current HEA is primarily attributed to precipitation hardening in addition to solid solution, grain boundary and dislocation hardening. The consolidated specimens undergo artificial ageing to study the ageing effect on the microstructure and mechanical properties of Al20Co20Cu20Zn20Ni20 HEA.
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ACCEPTED MANUSCRIPT For the peak aged samples, in addition to the aforementioned four terms, strengthening due to precipitate hardening contributes significantly. The formation of coherent precipitate of ordered FCC α phase (L12 structure) within disordered FCC (γ) matrix at 773 K as indicated in Figure 4 (a) significantly increases the hardness by impeding the motion of dislocation through the lattice. In shearing mechanism, various properties of the precipitates visʹ a visʹ that of matrix further decide the contribution from the following micro-mechanisms of
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deformation in a matrix containing coherent precipitates and contribute to strengthening [38]: (i) Coherency strain: attributed to strain field resulting from difference in lattice
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parameter of the matrix and the precipitate.
(ii) Stacking fault energy: arising due to significant difference in stacking fault energy of
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matrix and precipitate, affects glide of dislocation from matrix to precipitate as the fault width is a function of stacking fault energy.
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(iii) Ordered structure: very important in superalloys as precipitate has ordered structure, provides high temperature stability, strengthening attributed to creation of anti-phase
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boundary.
(iv) Modulus effect: precipitates with significantly different modulus than matrix affect the energy of dislocation while passing through them.
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(v) Interfacial energy and morphology: creation of additional surface due to cutting of
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precipitate by dislocations which results in an increase in surface energy. (vi) Lattice friction stress: different Peierls stress due to lattice friction in matrix and precipitate.
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Though it is difficult to estimate the contribution of individual strengthening mechanisms for precipitation hardening in the present HEA, it is expected that a combination of
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aforementioned strengthening mechanisms lead to three fold increase in hardness in the peak aged sample compared to the solutionized sample. With further ageing, at 873 K the hardness decreases due to the loss of coherency, increase in interparticle spacing and the dislocation (Orowan) looping around the particles [30]. Also, the decrease in hardness at higher ageing temperature for isochronal ageing and longer ageing time for isothermal ageing is attributed to the increases in size of the α phase particles [42, 43]. The results of isochronal ageing studies show that the peak hardness is obtained at significantly lower ageing temperature with finer distribution of α phase within the γ phase, but the required ageing time increases
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ACCEPTED MANUSCRIPT dramatically compared to the higher temperature and lower time required for reaching the peak hardness in sinter-aged specimen.
5. Conclusion The present investigation has demonstrated the effect of artificial ageing treatment on the
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investigated novel sinter-aged equiatomic Al-Co-Cu-Ni-Zn HEA. In particular, the following conclusions are drawn from the study:
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1. The sinter-aged Al20Co20Cu20Ni20Zn20 HEA at 1223 K can be solutionized at 1433 K for 96
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h followed by ice water quenching, to obtain single phase FCC structure. 2. Isochronal heat treatments of the solutionized specimens for 48 h at various temperatures
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lead to the formation of ordered FCC α (L12) phase within the grains of FCC γ phase. This contributed to increase in hardness and peak hardness was observed for the sample aged at
MA
773 K. after which the hardness decreased with further increase in ageing temperature. 3. Isothermal ageing at 773 K further corroborated precipitation hardening as there was an increase in hardness with ageing time with peak hardness for 48 hour ageing after which the
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hardness decreased.
4. The presence of nanostructured ordered FCC α (L12) precipitates in the FCC γ matrix can
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explain the ageing behaviour of Al20Co20Cu20Ni20Zn20 HEA .
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Precipitate size
temperature (K)
(nm)
773
12 ± 1
873
20 ± 2
973
69±8
1073
224 ± 25
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Ageing
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(Atomic %)
(Atomic %)
Al
14.7±0.22
4±0.24
Co
49.8±3.44
28.1±1.67
Ni
20.2±2.04
16.0±0.48
Cu
7.5±1.56
28.3±1.26
Zn
7.8±0.76
23.6±1.21
Total
100.00
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Table II: Composition of α precipitate and γ matrix measured using EDS attached to TEM
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ACCEPTED MANUSCRIPT List of Figures Figure 1: DSC heating and cooling cycles of the sintered equiatomic HEA. Figure 2: X-ray diffraction patterns of pellets aged at different ageing temperatures indicated in the figure. The patterns at the bottom are from pellet sintered at 1223 K and from the fully solutionized sample shown for comparison.
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Figure 3: BSE SEM micrographs obtained from sample: (a) fully solutionized ; aged at
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(b) 873K and (c) 1073K.
Figure 4: (a) TEM bright field micrograph and (c) corresponding SAD patterns obtained
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one β grain.
Figure 5: TEM micrographs of (a, b) specimen aged at 773 K, (c) low magnification, (d)
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List of Tables
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comparison in (a). (b) Hardness variation of the peak aged sample as a function of time.
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Table I: The precipitate size as a function of ageing temperature, measured using TEM bright field micrographs.
Table II: Composition of α precipitate and γ matrix measured using EDS attached to TEM
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1. Artificial ageing of the novel equiatomic Al-Co-Cu-Ni-Zn High Entropy Alloy 2. Isochronal heat treatments of the solutionized specimens for 48 h at various temperatures lead to the formation of ordered FCC α (L12) phase within the grains of FCC γ phase 3. Peak ageing was observed for the sample aged at 773 K 4. Isothermal ageing at 773 K further corroborated precipitation hardening
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