Sinter ageing of equiatomic Al20Co20Cu20Zn20Ni20 high entropy alloy via mechanical alloying

Materials Science & Engineering A 617 (2014) 211–218

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Sinter ageing of equiatomic Al20Co20Cu20Zn20Ni20 high entropy alloy via mechanical alloying Sutanuka Mohanty, N.P. Gurao, Krishanu Biswas n Department of Materials Science and Engineering, Indian Institute of Technology Kanpur, Kanpur 208016, India

art ic l e i nf o

a b s t r a c t

Article history: Received 3 July 2014 Received in revised form 16 August 2014 Accepted 19 August 2014 Available online 28 August 2014

The present investigation reports for the first time, the sinter ageing of equiatomic Al20Co20Cu20Ni20Zn20 high entropy alloy (HEA), being synthesized by high energy ball milling of elemental powder blend under protective argon atmosphere, followed by consolidation of the milled powder by spark plasma sintering at different temperatures and applied pressure of 50 MPa. The detailed X-ray diffraction and transmission electron microscopy (TEM) studies indicate the presence of single phase, FCC β supersaturated solid solution in the ball milled powder. However, the sintering of the as-milled powder reveals the formation of α with ordered FCC (L12) structure within the grains of FCC γ. The microstructural analysis using TEM shows the precipitation of near cuboidal shaped α phase within the grains of γ. The size and shape of the precipitates depend on the sintering temperature. Hardness measurement of the sintered alloys suggests age hardening of the as-milled powder during sintering. The sinter age hardening of HEA is attributed to the fine scale precipitation of α phase. Detailed variation of the hardness and microstructural evolution are reported here to elucidate this novel finding. & 2014 Elsevier B.V. All rights reserved.

Keywords: HEA Spark plasma sintering Precipitation hardening Hardness Multi-component alloys

1. Introduction The universe of alloys has been expanded by the addition of new class of metallic materials, known as high entropy alloys (HEAs), which is attracting vigorous research activities now-a-days [1]. By definition, HEA is an alloy consisting of at least 5 principal metallic elements in equal or near equal molar concentrations [1,2]. During the past decade, this new class of alloys has been discovered by Yeh et al. [2] and has closely been followed by many research groups in the world [3–10]. The presence of 5 or more principle elements in the HEAs has fundamental effects on the entropy of mixing, free energy change, phase selection as well as stability. Earlier, it was anticipated that multi-principle elements in alloys will lead to the formation of intermetallic phases, complex microstructure and poor mechanical properties. On the contrary, the microstructure of HEAs is found to be consisting of FCC and/or BCC phases. Therefore, these alloys have been extensively investigated from both scientific and technological viewpoints as they exhibit interesting fundamental physics and technological promise [11]. A brief literature survey indicates that the HEAs in bulk form can be prepared by two processing routes; rapid solidification and

n

Corresponding author. Tel.: þ 91 512 2596184; fax: þ91 512 2597550. E-mail address: [email protected] (K. Biswas).

http://dx.doi.org/10.1016/j.msea.2014.08.046 0921-5093/& 2014 Elsevier B.V. All rights reserved.

mechanical alloying followed by sintering. However, the HEAs, containing multiple elements with varying melting temperatures, prepared by the rapid solidification route, reveal segregation and inhomogeneous microstructure [3,12]. This limitation can only be circumvented by using a combination of mechanical alloying (MA) followed by consolidation using sintering. In addition, the MA renders the formation of microstructure even at room temperature with better homogeneity [7–10,13] by driving the system away from equilibrium. Thus, subsequent heat treatment during (or after) the MA process can lead to the generation of different microstructures. One such possibility is the precipitation of different phases from the supersaturated solid solution phases formed by the MA route. The precipitation of different precipitates can occur even during the heating and cooling cycle of the subsequent consolidation process and renders a possibility of sinter ageing [14]. Although the microstructure and properties of HEAs at room temperature have been discussed in most of the recent studies, only few investigations report the effect of ageing treatment, especially at high temperature [15–21]. In the case of Al0.5CoCrCuFeNi alloy, ageing treatment above 700 1C leads to the formation of AlNi BCC phase as well as AlCu-rich FCC nano precipitate, causing substantial increase in hardness [15]. Lee et al. [16] have reported ageing behavior in Al0.5CoCrNiTi0.5 HEA due to the transformation of a BCC phase to CrCo-rich σ phase in temperature range of 800–1000 1C. Tsao et al. [17] have investigated age hardening behavior of Al0.3CrFe1.5MnNi0.5 HEA and revealed that it is due to

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the formation of both AlNi and Cr5Fe6Mn8 precipitates within the Cr-rich BCC dendrites of the as-cast sample. The age hardening of CoCrFeNiMo0.85 HEA in the temperature range of 600–1000 1C, leads to precipitation of fine needles of (Mo.Cr)-rich σ phase in near equimolar FCC matrix. The σ phase can completely be transformed to Mo-rich μ phase during subsequent heat treatment at 900 1C, leading to decrease in hardness [18]. Lin et al. [19] have reported the ageing of Cu0.5CoCrFeNi HEA during heat treatment of as-cast alloy from 350 to 1350 1C. Although ageing does not lead to substantial change in hardness, it causes precipitation of Cr-rich FCC phase in the FCC matrix. In another case, maximum hardness has been obtained for Al0.3CoCrFeNi alloys with minor amount of C, Mo and Ti at 700 1C due to a host of precipitates; κ2 carbides, (Ni, Al, Ti)-rich phase and (Cr, Mo)-rich phase [20,21]. The present investigation reports the formation of novel microstructures during consolidation of equimolar single phase Al–Co–Cu–Zn–Ni HEA using the spark plasma sintering (SPS) process. It will be shown that the supersaturated solid solution formed by the MA process undergoes ageing during consolidation at different temperatures and thus leads to precipitation hardening. The detailed correlation between the hardness of the sintered specimens and the microstructural evolution will be described and discussed in the light of currently available literature. It is to be noted that the SPS is known as field assisted sintering technique (FAST), which allows obtaining of faster densification at lower sintering temperature as compared to the conventional sintering processes [14] due to the application of current (500–1000 A) during sintering. Thus, it leads to highly localized temperature increases and subsequent spark discharge between the powder particles, activating the particle surfaces by the removal of the oxide films, if present during sintering [14]. The present investigation, for the first time, reports the sinter ageing of equimolar Al–Co–Cu–Zn–Ni HEA.

2. Materials and methods Elemental powders of Al, Co, Cu, Ni, Zn (purity greater than 99.9%) were used as starting materials. Milling was carried out in Pulverisette-5 planetary ball mill (Fritsch, Germany) with tungsten carbide (WC) vials and balls. Powders were milled up to 20 h at a speed of 200 rpm with ball-to-powder ratio (BPR) of 20:1 in dry condition under the protective argon (Ar) atmosphere. Prior to the milling, the vials containing powder and balls had been evacuated to 10  3 Torr and were refilled with high purity (99.9%) argon gas. The milling operation had intermittently been stopped and high purity argon gas had been refilled to ensure the protective atmosphere was retained inside the vials. The samples were picked up from the vial after the selected interval of the milling time to characterize by X-ray diffraction (XRD) and scanning electron microscopy (SEM) and transmission electron microscopy (TEM). The spark plasma sintering (SPS) of the milled powders was carried out using a Dr. Sinter 515S apparatus (SPS Syntex Inc., Kanagawa, Japan) with a pulse on–off ratio of 12:2. The powder mixture, loaded in a graphite die of 0.015 m diameter, was placed inside the SPS chamber between two graphite electrodes under a vacuum level of 6  10  3 Torr. High purity Ar gas was then purged with a flow rate of 2 l/min in the chamber to ensure minimal oxidation of the powders. The powder was heated with the same heating rate of 100 1C/min to different sintering temperatures of 600 1C, 700 1C, 800 1C, 900 1C and 1000 1C. The powder mixture was held at the sintering temperature for 6 min before it was furnace cooled to room temperature. The temperature of 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 wall so that the recorded temperature difference

between the sample and the die was minimum. A uniaxial pressure of 50 MPa was applied throughout the sintering cycle. During the SPS experiment, while a constant voltage of 20 V was applied, the current flow varied around 900 Amp. Phase identification of the polished and flat sintered as well as milled powder samples was performed using Bruker, D8 FOCUS X-ray diffractometer (XRD) with V-filtered Cr Kα (λ ¼0.228976 nm) radiation. The sample was rotated at a speed of 15 rpm to prevent the effect of preferred orientation. The peaks in the diffraction patterns in each specimen were identified using International Committee for Diffraction Data (ICDD) database. The detailed analysis of the XRD patterns using Rietveld analysis allows calculation of the lattice parameters of different phases present in the MA and the sintered specimens. The phase distribution and compositional analysis of all the polished samples were carried out using a Carl Zeiss EVO 50, scanning electron microscope (SEM) coupled with an energy dispersive spectrometer (EDS, INCA PENTA FET X3), operated at 20 kV. Transmission electron microscopic (TEM) observations were carried out using a 200 kV (FEI Tecnai, 20UT) microscope. Samples for TEM observation were prepared by the standard preparation technique, which included cutting of 3-mm disks, grinding, dimpling, and subsequent ion-milling at 5 kV at an angle of 41 until perforation occurred. The density of all SPS processed samples was determined according to Archimedes' principle using distilled water. The theoretical density for the sintered composition was calculated following the rule of mixtures, considering the theoretical densities of pure elements. Vickers hardness of the sintered pellets was measured using a Vickers micro hardness tester (Bareiss Prüfgerätebau, GmbH, Germany) by applying 200 g load with a loading time of 10 s. In order to obtain a reliable measurement of hardness, the diagonals of Vickers indents were measured using SEM. At least ten measurements were carried out for 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 measurement on both sides of the SPS pellet for each pellet. The hardness of the MA powder was measured using nanotribometer (Hysitron make TI900) using a load of 200 μN and holding time of 10 s.

3. Results In the following, we shall describe the results for the milled powder and the sintered samples obtained during the investigation in details. 3.1. X-ray diffraction (XRD) analysis Fig. 1a shows the XRD patterns obtained from mechanically alloyed (MA) powder during the course of ball milling. The pattern obtained from the 0 h powder mixture (Al, Co, Cu, Zn and Ni) is also shown at the bottom for comparison. It is clear that high energy ball milling causes the formation of single phase FCC solid solution (β) after 15 h of ball milling. The peaks corresponding to β phase are broad, indicating nanocrystalline nature of the β grains. No other peaks could be detected to the limit of resolution of the XRD. Therefore, we have utilized the powder ball milled for 15 h for subsequent consolidation. Fig. 1b reveals the XRD patterns of the pellets consolidated using SPS at different sintering temperatures, which are indicated in the figure. Sintering leads to the formation of two new phases indicated by γ (FCC) and α (L12). Therefore, single phase β supersaturated solid solution transforms to γ and α phases during sintering. It is to be noted that sintering carried out at temperature lower than 500 1C does not yield sufficient density and thus the results obtained from these

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3.3. Transmission electron microscopic (TEM) analysis

β

∗ Zn

β

20 hrs

β

15 hrs

o Al • Cu ♦ Ni ∇ Co χ FCC β FCC

Peak Intensity (in a.u.)

β • βχ o ∇ ♦

50

10 hrs

βχ • ∗

o

5 hrs ♦

60

70

80

∗∗

0 hr

∗

∇

90

100

110

120

Angle 2θ (in Degree)

β FCC

Peak Intensity (in a.u.)

γ FCC

Solutionized Sample

α L12

SPS at 900 °C SPS at 800 °C

γ β

50

60

213

SPS at 700 °C

α α γ

SPS at 600 °C α

β

70 80 90 100 Angle 2θ (in Degree)

15 hrs

110

120

Fig. 1. (a) X-ray diffraction patterns of ball milled powder for different time durations of ball milling and (b) X-ray diffraction patterns of pellets sintered at different sintering temperatures indicated in the figure. The pattern at the bottom is from 15 h ball milled sample whereas the pattern at the top is for the fully solutionized sample.

samples are not reported here. For comparison, the XRD pattern of the fully solutionized sample (heat treated at 1160 1C for 24 h) is also shown at the top of Fig. 1b. The peaks of the sample are found to be at the same position as those for the MA powder.

3.2. Scanning electron microscopic (SEM) analysis The SEM micrographs of the as-milled powder (15 h) as well as the sintered samples are shown in Fig. 2. The secondary electron micrograph of the as-milled powder shows a combination of faceted and round particles of different sizes. The EDS analysis of the powders was carried out to measure the elemental composition. It indicates that the particles have near equiatomic composition. The SEM micrographs of the sintered samples are obtained in back scattered electron (BSE) imaging mode to identify the different phases present in the sample. The micrographs also indicate reduction in the fraction of voids with increase in the sintering temperature (Fig. 2b–e). In addition, one can observe the formation of a two-phase microstructure in the samples sintered at 800 1C (Fig. 2d) and 900 1C (Fig. 2e).

We carried out detailed TEM investigation of as-milled powder and all the sintered samples. However, we shall present here a few representative micrographs to bring out the salient results. Fig. 3a shows typical low magnification bright field micrograph of the MA powder (15 h), revealing faceted particles of different sizes. The inset shows a high magnification micrograph of one such particle with faceted morphology. It shows the agglomeration of a number of nano-sized grains in the particle. The selected area diffraction (SAD) pattern obtained from the nanoparticle is shown in Fig. 3a0 . The diffraction rings can consistently be indexed due to the β phase. No diffraction ring corresponding to α or γ can be detected. Careful high resolution microscopic analysis (not shown here) indicates that the MA powder particles are devoid of any precipitate. Fig. 4a shows a bright field micrograph of the specimen sintered at 700 1C using SPS. It reveals the presence of nano-sized γ grains with finer scale α precipitates within the grains. The SAD pattern obtained from such a microstructure (as shown in Fig. 4a0 ) indicates the presence of rings corresponding to γ (FCC) and α (L12) phases. Some of the α precipitates are marked by arrows in Fig. 4a. The detailed analysis of the dark field micrographs (shown as an inset of Fig. 4a) indicates that the size of the precipitates varies from 5 to 20 nm. Fig. 4b and 4b0 describe the results of the TEM investigation on Al20Co20Cu20Zn20Ni20 MA powder sintered at 800 1C. The low magnification bright field micrograph (Fig. 4b) reveals micron sized grains of γ phase containing a large number of well dispersed α precipitates. The higher magnification micrograph (inset in Fig. 4b) shows near cuboidal α precipitates having rounded corners. The composite SAD pattern obtained from both γ and α phases (Fig. 4b0 ) indicate that they bear a specific orientation relationship. The detailed analysis reveals the most prominent orientation relationship is given by {111}γ∣∣{111}α and [001]γ∣∣[001]α. A detailed analysis of micrographs allows us to generate a histogram for the size of the α precipitates (Fig. 4b00 ). It indicates that the precipitate size varies from 45 nm to 90 nm with majority of the precipitates being in the range of 65 to 80 nm. Fig. 4c shows the typical bright field micrograph of the specimen sintered at 900 1C, revealing uniform distribution of α precipitates within the γ grains. A higher magnification micrograph (inset of Fig. 4c) shows well grown α precipitates in one of the γ grains. The micro-diffraction patterns obtained from both the γ grain and a α precipitate are shown in Fig. 4c0 and c00 , respectively. The patterns can be indexed due to γ and α phases. Thus, it is clear that α precipitates nucleate from the supersaturated β solid solution changing the composition of the matrix, leading to the formation of the γ phase. We carried out detailed compositional measurement of the γ and α phases present in the sample sintered at 900 1C. Fig. 5 shows a high angle annular dark field (HAADF) micrograph of the specimen revealing α precipitate in the γ matrix. The elemental compositional analysis carried out by energy dispersive X-ray spectroscopy (EDS) on both the phases is indicated in Table 1. The error in measurements is also indicated as standard deviation. It can clearly be observed that precipitates are rich in Al, Ni and Co and the matrix is rich in Cu and Zn. Thus, there is definite partitioning of elements from the matrix and the precipitate. It is to be noted that the MA powder contains all the five elements in almost equal proportions. 3.4. Hardness measurement of the sintered samples Fig. 6 depicts the hardness data of the sintered samples. The hardness is plotted against the sintering temperature in Fig. 6a. It shows that hardness increases as a function of sintering temperature, reaching a maximum value of 649 78 VHN at 800 1C followed by a decrease of hardness for higher sintering

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Elements Al Co Ni Cu Zn Total

Atomic% 17.45±0.9 20.19±0.7 20.28±0.7 21.52±0.8 20.56±0.4 100.00

5 μm

5 μm

5 μm

5 μm

5 μm Fig. 2. SEM images of (a) ball milled powder (15 h) and pellets sintered at (b) 600 1C; (c) 700 1C; (d) 800 1C and (e) 900 1C. The inset in figure (a) shows the EDS measurement done on the MA powder.

Fig. 3. (a) TEM bright field micrographs and (a0 ) SAD pattern obtained from the as-milled powder.

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215

50 nm

[011]α ⎢⎢[011]γ 100 nm

b

Count

50 nm

14 12 10 8 6 4 2 0 45 50 55 60 65 70 75 80 85 90 Precipitate Size (nm)

α

γ [112] γ

[001]

α

Fig. 4. TEM micrographs and corresponding SAD patterns obtained from specimens sintered at (a) 700 1C; (b) 800 1C and (c) 900 1C. The insets in figures (a), (b) and (c) show higher magnification dark field and bright field micrographs revealing the precipitates with (a0 ), (b0 ), (c0 ) and (c0 0 ) showing corresponding SAD patterns whereas (b0 0 ) shows a histogram for the size of the precipitates for the specimen sintered at 800 1C.

temperatures. This is a typical characteristic of age hardening in metallic systems [22]. The hardness of the as-milled powder is also shown for comparison. The hardness of the ball milled powder has

been measured by using a load of 200 μN and Vickers indenter in a nanotribometer. The hardness of the MA powder is found to be 210 720VHN. Fig. 6b shows the hardness variation as a function of

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50 nm

100 nm

Fig. 5. HAADF image of the pellet sintered at 900 1C with inset showing a higher magnification micrograph. The compositions of α and γ matrix phases are indicated in Table 1.

Table 1 Composition of α and γ phases measured using EDS attached to TEM. Elements

Precipitate (α) Atomic(%)

Matrix (γ) Atomic (%)

Al Co Cu Zn Ni

15.7 7 0.16 48.87 3.44 6.17 1.54 8.17 0.93 21.2 7 1.03

3 7 0.85 27.1 7 1.19 27.3 7 1.44 23.6 7 1.34 18.9 7 0.43

the ageing time of the MA powder sintered at 800 1C. This is obtained by varying the cooling rate in the SPS run after a fixed holding period (6 min) at 800 1C. The time varies from 5 min to 70 min. The result reveals that the hardness follows similar characteristics as in Fig. 6a. The typical SEM images of the indent on the sample are also shown as inset in each figure. For the fully solutionized specimen (as indicated in the XRD pattern Fig. 1b), the measured hardness is found to be 160 730 VHN. Therefore, sintering of the MA powder leads to three-fold increase in hardness, which is mainly due to the formation of α precipitates. 4. Discussion The present investigation has clearly demonstrated the sinter ageing of novel Al-containing multicomponent equimolar Al–Co– Cu–Zn–Ni alloy. The microstructural analysis conclusively shows the precipitation of ordered α phase from the supersaturated β solid solution. The hardness measurements indicate the age hardening of the MA powder during consolidation, which is a novel finding. However, it is important to discuss the results in the light of available literature. Two aspects require a detailed explanation. These are: (A) the formation and stability of nanocrystalline HEA in equimolar Al–Co–Cu–Zn–Ni alloy system and (B) sinter ageing behavior during consolidation of the MA powder. 4.1. Formation and stability of HEA in equimolar Al–Co–Cu–Zn–Ni system via mechanical alloying The detailed XRD as well as TEM investigations of the MA powder indicate that it is possible to form single phase HEA (β) in the multicomponent equimolar Al20Co20Cu20Zn20Ni20 alloy system. The XRD patterns of the MA powder clearly demonstrate that β phase forms after 15 h of ball milling under protective

Fig. 6. Hardness plot showing (a) Vickers hardness as a function of sintering temperature and (b) Vickers hardness as a function of time for the specimen with the highest hardness. The hardness of the as-milled powder (15 h) and the fully solutionized specimens are also shown for comparison.

atmosphere. This is a unique HEA phase-forming alloy system since it has not been reported in the literature so far. We shall now make an attempt to discuss the formation and stability of HEA from thermodynamical view point. The multiprinciple element HEAs are reported to be stabilized by high configurational entropy of mixing (ΔSconf) [2], given by n

ΔSconf ¼  R ∑ ðxi ln xi Þ i¼1

ð1Þ

where R is the universal gas constant and xi is the mole fraction of the component i present in the alloy. The ΔSconf ¼1.61 R for i¼5 with equimolar concentration. This value of ΔSconf is reported to be higher than the entropy of mixing of many of the metals and intermetallic phases and thus it leads to substantial lowering of free energy of mixing (ΔGmix ¼ ΔHmix  TΔSconf) of the solid solution phases, since ΔHmix is low for most of the metallic systems [23]. Thus, it causes the formation of stable solid solution phases. However, we also need to consider other factors responsible for stabilizing the random solid solution phases in these multicomponent alloys. Four additional factors have been reported to be considered to explain the stability of HEAs: (i) thermodynamic parameter (Ω Z1.1); (ii) atomic size difference (δ r6.6%); (iii) difference in Pauli's electronegativity (Δχ) and (iv) valence electron concentration (ψ). The last three factors are

S. Mohanty et al. / Materials Science & Engineering A 617 (2014) 211–218

known as Hume–Rothery factors. Mathematically, thermodynamic parameter, Ω [23], is given by

Ω¼

T m ΔSconf jΔH mix j

ð2Þ

where T m ¼ ∑ni¼ 1 C i ðT m Þi is the melting temperature of the single phase multicomponent alloy and ΔH mix ¼ ∑ni;j ¼ 1 Φij C i C j is known as the enthalpy of mixing for multi-component alloy system with n mix elements. For simplicity, we use Φij ¼ 4ΔH mix is ij , where ΔH ij the enthalpy of mixing of the equimolar binary alloy with i and j atoms [24]. It is to be noted that Ω is obtained by combining
ΔHmix and ΔSconf. For Ω Z1.1., T m ΔSconf =
ΔHmix j Z 1:1, i.e., TmΔSconf Z1.1∣ΔHmix∣ and thus, entropic contribution is higher than enthalpic contribution. The parameter δ is also given by [23] sffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi  n r 2 δ ¼ ∑ Ci 1  i ð3Þ r i¼1 where Ci is the concentration of ith component in the alloy and r is the average atomic radius. The size factor as well as ΔHmix for the binary equiatomic alloy composition in Al–Co–Cu–Zn–Ni system is shown in Table 2. On the basis of Miedema macroscopic model for binary liquid alloys the ΔHmix has been obtained [25]. The estimated values of all the relevant parameters for the binary couples and the equiatomic Al20Co20Cu20Zn20Ni20 alloy are shown in Tables 2 and 3, respectively. The values of Ω and δ clearly indicate that the HEA will be stable at room temperature for the multicomponent alloy. One also needs to consider two other factors responsible for solid solution formation. Guo et al. [26,27] have shown that the differences in electronegativity (Δχ) do not have much effect on the formation of solid solution phase in HEA forming systems. However, it has been reported that the valence electron concentration (ψ) can be used to predict the formation of stable FCC or BCC phases in these alloy systems. It has been shown that (a) FCC phase can only exist for ψ Z8.0, (b) FCC and BCC phases can co-exist for 6.87 r ψ o8.0 and (c) only BCC phase can exist for ψ o6.87. The calculated values of Δχ and ψ are also shown in Table 2. It can be seen that for the investigated alloy, only FCC phase will exist in the MA powder. Our experimental Table 2 The size factor and the enthalpy mixing for various binary equiatomic (or equimolar) alloys in the multicomponent Al–Co–Cu–Zn–Ni HEA [25]. Binary alloy (A–B)

Al–Co Al–Cu Al–Zn Al–Ni Co–Cu Co–Zn Co–Ni Cu–Zn Cu–Ni Zn–Ni

Size factor A–B

B–A

14.47 12.05 2.65 14.93 2.11 10.32 0.4 8.39 2.57 11.96

12.64 10.75 2.58 12.99 2.16 11.51 0.39 9.15 2.50 10.68

217

investigation also shows a similar result. Thus, it is possible to explain the formation and stability of HEA phase in the equiatomic quinary alloy at room temperature, as the multiprinciple element HEAs contain supersaturated solid solution, which upon further heat treatment can undergo transformation.

4.2. Sinter ageing behavior during consolidation of the MA powder The present investigation reveals that it is possible to adopt sinter ageing during spark plasma sintering of Al20Co20Cu20Zn20Ni20 HEA to obtain sintered compact containing L12 type α precipitates within the FCC γ matrix grains. The detailed hardness measurements conclusively prove that typical precipitation hardening takes place during the consolidation of the powder. To the best of the authors’ knowledge, this is the first report of sinter ageing of the complex multicomponent HEA. The detailed microstructural analysis indicates the formation of fine scale near cuboidal shaped nano-sized α precipitates with ordered L12 structure in the micron-sized FCC γ grains depending on the sintering temperatures. In the following paragraphs we shall discuss about the sinter ageing process and the microstructural evolution. The mechanically alloyed powder reveals the presence of single FCC solid solution (β) having lattice parameter aβ ¼0.363 nm. During subsequent sintering experiments, this supersaturated solid solution becomes unstable and transforms to another FCC solid solution (γ) and precipitates (α). The peak positions of the phases in the XRD patterns (see Fig. 1b) indicate a small difference in the lattice parameters of the supersaturated β and γ solid solution phases. The lattice parameter of the γ (aγ) sintered at different sintering temperatures determined using Rietveld analysis XRD patterns is listed in Table 4. Its value is found to vary between 0.366 and 0.367 nm. Therefore, there is very small β change in the lattice parameters of the FCC phases (Δaγ ¼ 0:004 nm or 1.1%) during sinter ageing of the MA powder. Such a change in lattice parameter can be correlated to chemical partitioning of solute elements during the transformation of β to γ and α. Thus, we term β as supersaturated solid solution formed during mechanical alloying with γ being the solid solution. This is similar to the γ0 (Ni3Al) phase in Ni-based superalloys that precipitates in the γ phase which is a solid solution of Ni with other elements. The presence of high volume fraction of ordered precipitate provides strength to the superalloys at an elevated temperature

Enthalpy of mixing (ΔHmix) in KJ/mol Table 4 Lattice parameter of precipitate (aα) and matrix (aγ) determined using XRD. Precipitate size as a function of sintering temperature is listed in the last column. The precipitate size for 600 1C is obtained from the XRD data.

 19 1 1  22 6 5 0 1 4 5

Sintering temperature (1C)

Precipitate (aα) (nm)

Matrix (aγ) (nm)

Precipitate size (nm)

600 700 800 900

0.3567 0.08 0.3577 0.08 0.3567 0.07 0.3587 0.09

0.3667 0.06 0.3677 0.09 0.3667 0.03 0.3677 0.03

15 76.7 18 76.1 72 78.5 158 718.5

Table 3 The thermodynamic parameters such as δ, (ΔHmix), (ΔSmix), Tm, ΔG, Ω, VEC and Δχ calculated for the alloy system under study. Atomic size difference δ (%)

Enthalpy of mixing ΔHmix (KJ/mol)

Entropy of mixing ΔSmix Melting (JK  1 mol  1) temperature Tm (K)

ΔG Thermo-dynamic (J mol  1) parameter Ω

Valence electron compound VEC

Electro-negativity differenceΔχ

5.9

 7.04

13.38

 11,054

9

0.29

1295.98

2.46

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[28]. This is mainly due to coherency strain between γ0 precipitate and γ matrix in the superalloys. In the present case, a similar observation is made for α precipitates within γ grains. Table 4 also lists the lattice parameter of α (aα) phase measured using Reitveld analysis of the XRD patterns as a function of sintering temperature. aαvaries from   0.356 to  0.357 nm. Thus, the misfit (defined as 2 ðaα  aγ Þ = ðaα þ aγ Þ ) is significantly γ large (Δaα ¼ 2:8%). In classical Rene88DT alloy, the lattice parameters of the FCC phase and L12 ordered phases are 0.359 and γ 0.358 nm and thus the misfit is smaller (Δaγ 0 ¼ 1:1%) [29,30]. This has two effects on the microstructure and hardness of the specimens. The large misfit will lead to substantial coherency strain between γ and α phases. In fact, it has been indicated that the large misfit can lead to loss of coherency. However, present investigation clearly suggests that the precipitates and the matrix bear an orientation relationship indicating that coherency is maintained. It is to be remembered here that there is a relaxation from the stringent lattice parameter mismatch in the Hume– Rothery rule from 15% for two elements to 6.6% for multi component HEA systems [27,31]. This leads to a highly distorted lattice for the HEA and any phase formation occurring with partitioning of elements will have another distorted lattice. Therefore, the lattice parameter mismatch between the precipitate and the matrix phase in HEAs can be expected to be higher than that of conventional age hardening alloys. The precipitate size in the γ grains has been measured using TEM micrographs. The average precipitate size as a function of sintering temperature is listed in Table 4. Correlating hardness with precipitate size indicates the coherency and is assumed to be maintained until a critical size of 7075 nm of the precipitate in the γ matrix. Therefore, peak hardness can be achieved for the sample sintered at 800 1C. The hardness value shows a downward trend at 900 1C, which can be related to the loss of coherency due to the growth of the precipitates. The precipitation mechanism leading to the formation of α from β is under study and will be reported separately. However, it definitely involves diffusion of elements in the multicomponent alloy and therefore will come under diffusional transformation. The entire process of precipitation is modified due to the sluggishness of diffusion of different species in HEAs compared to simple substitutional solid solutions [32]. However, the net effect of precipitates in terms of blocking the path of dislocations is similar to that in any precipitation hardening alloys, and a similar ageing behavior is observed [22]. There is a three-fold increase in hardness for the peak aged sample compared to that of as-milled powder and the solutionized sample (Fig. 5a.) for the HEA, which is comparable to the optimum peak hardened aluminum alloys. Various parameters like coherency strain, chemical environment, ordering, stacking fault and difference in modulus contribute to the extent of precipitation hardening in age hardenable alloys. In the present case, it is expected that HEA will have higher contribution from chemical hardening and ordering of the precipitate. Thus, the peak in hardness followed by a gradual decrease due to the loss in coherency is observed for HEA similar to conventional precipitation hardening alloys [33]. The presence of the ordered L12 α phase in FCC γ matrix in this HEA is similar to that in nickel based superalloys. However, in this case there are some major differences in the precipitate and matrix characteristic. Unlike superalloys, the matrix phase is nanocrystalline and the precipitate fraction is much lower (15– 20% versus 70% in superalloys) than that observed in nickel based superalloys [30,31]. HEA are expected to be stable at high temperatures (up to 950 1C) and the precipitates are likely to behave like the ones in nickel based superalloys. The sluggish diffusion in HEA can reduce the extent of diffusion creep at high temperatures. This HEA may even be immune to rafting

(progressive elongation of cuboidal precipitates in the presence of stress and temperature) which is present in superalloys [34]. Therefore, it may be possible to achieve the best of both worlds in terms of obtaining high temperature stability like superalloys without the problem of precipitation coarsening or rafting. The detailed study on the stability of the precipitates is also under investigation and will be reported separately.

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