Comparative study of microstructure and mechanical properties of thermo-mechanically processed Al_{(0.2, 0.5)} CoCrFeNiMo_0.5 high-entropy alloys

ABSTRACT

The main objective of the present work is to produce Al_xCoCrFeNiMo_0.5 alloy through a wrought process. Al_xCoCrFeNiMo_0.5 high-entropy alloys with two different Al contents (x = 0.2 and 0.5) were prepared by vacuum arc melting followed by thermo-mechanical processing. Microstructural evolution was investigated after hot forging and subsequent homogenisation using XRD, SEM, and TEM. The final microstructure corresponding to Al_0.2 and Al_0.5 are BCT (sigma) in the FCC matrix and BCT (sigma) and ordered BCC together in the FCC matrix respectively. Room-temperature mechanical properties were investigated using uniaxial compression and hardness measurements. There is an apparent increase in strength and a decrease in ductility for the Al_0.5 alloy. However, a higher ductility was observed in Al_0.2 alloy with a compromised reduction in strength.

KEYWORDS:

• High-entropy alloy
• microstructure
• forging
• deformation
• fracture

1. Introduction

The new alloying strategy of mixing multiple principal elements in large quantities, in the range of 5-35%, is drawing much scientific curiosity [1,2]. This new class of materials, called high-entropy alloys, have potential applications as structural materials because of their high strength and high plasticity [3–8], high hardness [9], good thermal stability [10,11], high wear resistance [12], and high irradiation resistance [13]. Furthermore, recent studies suggest that these alloys can be potential materials for high-temperature applications.

AlCoCrFeNiMo is one such multi-component system with a gradual change in the microstructure from FCC + BCT to BCC + BCT and finally to BCC1 + BCC2 with increasing Al content [14]. Although available literature on AlCoCrFeNiMo alloys suggests that they can be good candidates for high-temperature applications [15], most of the studies on this alloy are confined to either as-cast or following subsequent heat treatment [11,15–21].

Thermo-mechanical processing is a combination of thermal treatment followed by mechanical working such as forging and rolling. This is not only used to process the primary cast ingot to final shapes but also to enhance mechanical properties by tailoring the microstructure. Furthermore, the method breaks the as-cast dendritic type coarse microstructure and reduces macro-segregation, thereby further improving the mechanical properties compared to as-cast components. No studies are available on the effect of thermo-mechanical processing on the microstructure and mechanical properties of Al_xCoCrFeNiMo_0.5 high-entropy alloys. The present work investigatese the effect of the thermo-mechanical processing route on the microstructure and mechanical properties of Al_0.2 and Al_0.5 high-entropy alloys (HEAs).

2. Experimental details

Two non-equiatomic HEAs, Al_0.2CoCrFeNiMo_0.5 and Al_0.5CoCrFeNiMo_0.5, were prepared through vacuum arc melting of pure elements (purity 99.9%) in the form of pellets, with 20 grams in each coupon. Alloy coupons were re-melted five times in an argon atmosphere for better chemical homogeneity and subsequently hot forged (HF) at a temperature of 900 °C and then homogenised (HM) at 1200 °C for 6 h. Specimens for microstructural examination in a scanning electron microscope (model: FE-SEM JEOL-JSM 7800F; JEOL, Japan) were polished using SiC grit papers. The samples were electro-polished using 90% methanol and 10% perchloric acid solution using the following parameters: voltage 20 V, time 10 s, and temperature −15 °C. Individual phases and corresponding crystal structures present in both alloys were identified using a transmission electron microscope (model: JEM-2100; JEOL, Tokyo, Japan). The samples used for TEM studies were polished to 80 μm in thickness, followed by twin-jet polishing in an electrolyte of 90% methanol and 10% perchloric acid (20 V at −15 °C). Micro-Vickers hardness (model: Durascan; Emco Test, Kuchl, Austria) at a load of 5 kgf for 10 s dwell time were used to measure the hardness of the processed specimens. Compression experiments were performed on the forged and homogenised samples with a rectangular cross-section of l/d ratio 1.5, with an initial strain rate of 1 × 10⁻³ s⁻¹ (model: Instron 5967, Instron Inc., USA).

3. Results and discussion

X-ray diffraction patterns of Al_0.2 and Al_0.5 after hot forging (HF) and subsequent homogenisation at 1200 °C for 6 h (HM) are shown in Figure 1a. It can be observed that the Al_0.2 and Al_0.5 alloys consist of FCC + BCT (sigma) and FCC + BCC + BCT (sigma) phases respectively. Microstructures captured using the scanning electron microscope in backscattered mode correspond to the Al_0.2 and Al_0.5 alloys in HF and after HM shown in Figure 2a–d. The dendritic structure is still prominent in the case of Al_0.2 compared to Al_0.5, with the former being entirely transformed to an equiaxed grain structure during subsequent homogenisation. This may be the reason for significant changes in the peak intensities in the XRD pattern of Al_0.2 before and after HM treatments as compared to Al_0.5 (Figure 1).

Figure 1. (a) XRD of Al_0.2, Al_0.5 in (HF) and (HM) conditions showing FCC + BCT phase in Al_0.2 and FCC + BCT and BCC Al_0.5, EBSD maps of (b) Al_0.2 HM and (c) Al_0.5 HM.

Figure 2. Backscattered electron microstructure of (a) Al_0.2 HF, (b) Al_0.2 HM, (c) Al_0.5 HF, (d) Al_0.5 HM, TEM selective-area diffraction patters of different phases present in (e) Al_0.2 HM, (f) Al_0.5 HM and elemental distribution in (g) Al_0.2 HM and (h) Al_0.5 HM.

TEM micrographs and diffraction patterns of the individual phases present in the Al_0.2 and Al_0.5 alloys are shown in Figure 2e and f, respectively. It is clear from Figure 2e that the matrix phase and the secondary phase in Al_0.2 corresponds to FCC and BCT (sigma) phase. Similarly, the diffraction patterns of the FCC matrix phase and the secondary BCT phase are inconsistent with the XRD of Al_0.5. However, the diffraction pattern corresponding to the third BCC phase in Al_0.5 consists of satellite spots, suggesting that the third phase, the indexed BCC phase in the XRD, corresponds to Al_0.5 (Figure 1a) and is rather an ordered BCC (B2) structure. Elemental distributions after homogenisation treatment of the Al_0.2 and Al_0.5 HEAs are shown in figure 2g and h. In addition, the compositions (in atomic %) of the individual phases in each of the alloys are listed in Table 1. The overall compositions of the alloy determined using SEM-EDS and ICP-OES (inductively coupled plasma – optical emission spectrometry) along with the targeted nominal composition during melting are also included in Table 1. It is suggested that the bright sigma phase in both Al_0.2 and Al_0.5 is rich in Cr and Mo and the FCC matrix phase in both alloys consists of Co, Cr, Fe, and Ni in nearly equiatomic proportion with minor amounts of Al and Mo. The gray colour phase in Al_0.5 alloy, which is an ordered B2 structure, is rich in Ni and Al, suggesting that it might be a NiAl-based ordered phase [21].

Table 1. Chemical composition of individual phases present in each alloy determined using SEM-EDS along with the bulk composition calculated using ICP OS.

Alloy	Method used	Al (at %)	Co (at %)	Cr (at %)	Fe (at %)	Ni (at %)	Mo (at %)
Al0.2 1200°C 6hr	Nominal	4.3	21.3	21.3	21.3	21.3	10.6
Total	ICP OS	4.6	21.9	21.2	21.9	20.4	10
Total	SEM-EDS	4.8	21.7	22.2	22.5	20.2	8.6
FCC	Dark	5.4	22.3	20.6	23.3	22.2	6.1
BCT	Bright	0	19.7	29.3	20	12.1	19
Al0.5 1200°C 6hr	Nominal	10.0	20	20	20	20	10
Total	ICP OS	11.2	19.5	21.9	19.5	19.2	8.9
Total	SEM-EDS	10.8	19.2	22.7	20	19.4	7.9
FCC	Dark	7.3	22.5	20.5	24.8	21.6	3.4
BCC(B2)	Light dark	30.2	14.2	6.7	12	35.4	1.6
BCT	Bright	3.2	19.2	32.4	20.8	9.6	14.8

The deformability of the matrix plays a significant role during the thermo-mechanical processing of any alloy to its final shape with minimum defect density such as cracking. The matrix in both HEAs is FCC which is softer and ductile compared to either the sigma phase or ordered BCC [19]. It is clear from the microstructural analysis that it is a continuous phase in Al_0.2 before and after thermo-mechanical processing. In contrast to Al_0.2, though the matrix in Al_0.5 is continuous in the as-cast condition (Figure S1(a) and S1(b)), it breaks down to an elongated discontinuous phase during forging and further to an equiaxed entirely discontinuous phase after homogenisation. Owing to this discontinuous matrix, though it is possible to forge both HEAs in the as-cast form, it is comparatively difficult to roll Al_0.5 at 900 °C without forming edge cracks.

The values of the micro-Vickers hardness of the Al_0.2 and Al_0.5 alloys measured after hot forging and subsequent homogenisation are shown in Figure 3a. Although the homogenisation treatment has no influenceon the hardness, it is clear from Figure 3a that the hardness increases with increasing Al content. The room-temperature compressive stress-strain flow behaviour corresponding to Al_0.2 and Al_0.5 tested in homogenised conditions is shown in Figure 3b. From stress-strain curves, the yields stress, the ultimate tensile stress, and the percentage elongation for Al_0.2 are 1007, 1683 MPa, and 14.5%, respectively, while for Al_0.5 they are 986, 2086MPa, and 6.7%, respectively. It is observed that Al has a strong influence on the flow behaviour of Al_xCoCrFeNiMo_0.5 HEA alloy. There is also a significant increase in the strength as well as a decrease in ductility with increasing Al content. The high yield strength in Al_0.5 may arise from the presence of the ordered BCC (B2) phase along with the hard, brittle sigma phase. Furthermore, the volume fractions of the hard phase in Al_0.2 (sigma phase) and Al_0.5 (sigma phase + B2) are 17 and 39%, respectively. The grain sizes and fractions of individual phases are included in Table 2, suggesting that the presence of a large volume fraction of the hard phase in Al_0.5 increases significantly the load-bearing capacity of the alloy. On the other hand, increasing the hard phase in Al_0.5 interrupts the continuity of the soft FCC matrix, especially after homogenisation, thereby making it difficult to maintain the strain continuity, causing stress concentrations at the boundary-matrix interface. This results in early crack initiation and subsequent propagation as shown in Figure 3c, causing a decrease in ductility.

Figure 3. (a) Micro-Vickers hardness of Al_0.2 and Al_0.5 after hot forging (HF) and after homogenisation (HM) at 1200°C for 6 hr (b) Compression flow behaviour of Al_0.2 and Al_0.5 in homogenisation condition. (c) the fracture surface of the Al_0.5 in HM condition.

Table 2. Phase fraction and grain size of the individual phases present in Al_0.2 and Al_0.5 HEAs.

		Phase fraction	Grain size (µm)
Al0.2 HF	FCC	80	–
	BCT	20	–
Al0.2 HM	FCC	77	–
	BCT	23	–
Al0.5 HF	FCC	32	3.2 ± 1.1
	BCT	47	3 ± 0.9
	B2	21	2.2 ± 0.6
Al0.5 HM	FCC	32	3.4 ± 1
	BCT	49	2.8 ± 0.9
	B2	19	2 ± 0.7

4. Conclusions

Wrought high-entropy alloys of Al_(0.2&0.5)CoCrFeNiMo_0.5 were successfully produced through a thermo-mechanical processing route and the resulting microstructure and mechanical properties studied. Transformation in the microstructure from BCT sigma phase in the FCC matrix to BCT sigma and ordered BCC (B2) phases together in the FCC matrix was observed with increasing Al concentration. The soft FCC matrix is continuous in the Al_0.2 alloy after forging even after homogenisation treatment. In contrast, the elongated discontinuous phase in the case of the Al_0.5 alloy during forging which thereafter transformed to a fully equiaxed discontinuous phase after homogenisation, making it difficult to accommodate strain during subsequent processing operations such as rolling. The maximum strengths to fracture and the fracture strain corresponding to the Al_0.2 and Al_0.5 alloys are 1683 and 2086MPa, 6.7, and 14.5%, respectively. An increase in strength with increasing Al arsies from an increase in the volume fraction of the hard phase with Al. A decreasing ductility in the Al_0.5 alloy arises from a difficulty in achieving strain continuity because of the presence of a discontinuous matrix in the Al_0.5 alloy compared to the Al_0.2 alloy_.

Acknowledgments

The authors would like to acknowledge the financial support of the DST-FIST (grant no. SR/FST/ETI-421/2016) program for funding the FESEM-EBSD facility and Science and Engineering Research Board, Government of India (grant no. ECR/2017/001278) for financing the project. The authors also like to acknowledge Dr. Mithun Palit for helping in the alloy preparation.

Disclosure statement

No potential conflict of interest was reported by the author(s).

Additional information

Funding

This work was supported by SERB-ECR: [Grant Number ECR/2017/001278]; DST-FIST: [Grant Number SR/FST/ETI-421/2016].