Effect of multiple oxides on the mechanical properties of CoCrFeMnNi high-entropy alloy matrix composites

Introduction

High-entropy alloys (HEAs) are multi-component alloys composed of more than five main elements of 5–35 at.-% [1]. HEAs generally exhibit outstanding hardness and strength due to solid solution strengthening by significant distortion of lattices [2–5]. Recently, efforts have been made to further increase the strength of HEAs by producing fine and uniform microstructures using powder metallurgical routes [6–8] and by generating ex situ or in situ hard phases, such as nano-scale oxides [9,10], carbides [11,12], and reinforcement [13,14] in the HEA matrix.

Among HEAs, the CoCrFeMnNi alloy is a remarkable HEA because of its excellent mechanical properties at cryogenic temperatures and high phase stability over a wide temperature range [15,16]. However, the relatively low yield strength at room temperature should be overcome to enable its application in various fields. Many studies have been conducted to improve the mechanical properties of CoCrFeMnNi HEAs. For instance, the use of additional alloy elements (i.e. Al, V, Nb, Ti) in the second phase of the HEAs can improve its strength [17,18]. In addition, the strength of the HEAs can be also increased by grain refinement by severe plastic deformation [19]. Furthermore, CoCrFeMnNi-matrix composites reinforced with nanoparticles have outstanding mechanical properties. TiC- and SiC-reinforced CoCrFeMnNi HEA composites have been manufactured using powder metallurgy, and they showed fine grain size, high yield strength, and high strain hardening coefficient [11,12].

Oxide-dispersion strengthening (ODS) is a potential strategy to enhance the strength of HEAs, as it inhibits the dislocation motion by oxides. The oxides restrict the grain growth via the pinning effect [20,21]. Oxide-dispersion-strengthened HEA matrix composites have been widely studied. For example, an Al_0.3CoCrFeMnNi-based HEA composite was fabricated with 0–3 vol.-% Y₂O₃ [9]. Another study investigated the synergic effect of nano-sized titanium oxide dispersion and coherent precipitates in the intermetallic phase produced by powder metallurgy [10]. However, most studies on HEA-based composites focus on the development of new materials with exotic properties. However, various types of oxides newly formed during processes and their effects on mechanical strength have not been sufficiently investigated.

In this study, we investigated the effect of milling and heat treatment durations on the in situ formation of second phases and various types of oxides as well as on the grain size of HEAs and dispersion of ex situ Y₂O₃ nanoparticles. Furthermore, their contribution to the mechanical properties of ODS-HEA matrix composites is discussed.

Materials and methods

CoCrFeMnNi-based ODS composites were produced by hot-pressing a mechanically milled powder. Water-atomised CoCrFeMnNi equiatomic alloy powder (particle size <150 μm) and nano-sized Y₂O₃ particles (particle size: 50–80 nm) were mixed via high-energy ball-milling using an attrition mill (KMC-2BV, KMC Co. Ltd.). For that, 97 g of CoCrFeMnNi HEA powder and 3 g (3 wt-%) of Y₂O₃ particles were placed in a stainless steel chamber with 1000 g of milling media, namely stainless steel balls (5 mm diameter). The milling was conducted at 500 rev min⁻¹ for 12, 24, and 48 h under an Ar atmosphere with 3 wt-% process control agent, namely stearic acid (CH₃(CH₂)₁₆CO₂H, Sigma Aldrich Korea Co, Ltd.). After the milling process, the stearic acid was removed by heating at 500 °C for 20 min. During the milling, cold water circulation was used to prevent excessive heat generation. Subsequently, the ball-milled powder was consolidated in a stainless steel mould by hot-pressing at 600°C with an applied pressure of ∼70 MPa for 2 h. Hot-pressed pallets were further heat-treated at 800°C for 12 h under an Ar atmosphere for all pallets.

The morphologies of the starting powder and ball-milled powder were observed by scanning electron microscopy (SEM, JEOL JSM 2001F). The phases of the powder and hot-pressed pallets were investigated using X-ray diffraction (XRD, Rigaku Ultima iii X-ray diffractometer and SmartLab) with Cu Kα radiation. The samples were scanned at a 2θ scan range of 20°–110° with a step size of 0.02° and a scan speed of 10 °/min. X-ray photoelectron spectroscopy (XPS, Thermo, K-alpha) was used to identify the type of oxides in the heat-treated ODS composites. The surface of the samples was etched with an Ar ion at 500 eV for 200 s to remove the oxide layer. The microstructure of the heat-treated samples was investigated using high-resolution transmission electron microscopy (HR-TEM, JEOL JEM-2100F). Specimens for the HR-TEM analysis were prepared using a focused-ion-beam (FEI Nova Nanolab).

The mechanical properties of the ODS composites were measured in a compressive test using a universal testing machine (UTM, R&B, RB Model 301 Unitech M) with a 5 kg load cell. Compressive tests were conducted at a strain rate of 1 × 10⁻⁴ s⁻¹. The specimens for the compressive test were prepared in a 3 mm × 3 mm × 3 mm in width, length, and height. Two tungsten carbide plates were used to sandwich the specimen, and a sprayed film of boron nitride was used as a lubricant to minimise the effects of friction.

Results and discussion

Figure 1 shows SEM images of the starting CoCrFeMnNi (Figure 1(a,b)) and Y₂O₃ powder (Figure 1(c,d)). The alloy powder was spherical, but satellites were observed on the surface because the surface roughness of the powder increased with increasing powder size during the water atomisation process. The cooling rate of the capacity decreased as the heat transfer coefficient decreased with increasing powder size [22,23]. Therefore, relatively small-sized powders condensed on the surface of large powders during atomisation. As the average particle size of the Y₂O₃ was 50–80 nm, they agglomerated because of their large specific surface area.

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Figure 1

SEM images of (a,b) starting HEA particles, and (c,d) nano-size Y₂O₃ particles.

Figure 2 shows the SEM images of the ODS composite powders after high-energy ball-milling for various milling times (12, 24, and 48 h). The morphology of the ODS composite particles changed from spherical to flaked during the high-energy ball-milling. The average particle size of the composites decreased as the milling time increased from 12 to 48 h. The ball-milling process involved repeated fracturing and cold welding. As seen in the magnified images (Figure 2(b,d,c)), fractured small powders were cold welded on the surface of relatively large powders. In addition, the elemental maps (Supplementary Figure 1) revealed that the Y₂O₃ nanoparticles were more uniformly dispersed in the HEA powder as the milling time increased. During mechanical milling, relatively soft HEA powder can plastically deform owing to the high mechanical energy, and hard Y₂O₃ particles would be embedded into the HEA powder and gradually be dispersed through the powder mass flow [24].

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Figure 2

SEM images of the ODS-HEA composite powders after milling for various times: (a,b) 12 h, (c,d) 24 h, and (e,f) 48 h.

Figure 3 shows the XRD patterns of the ODS composites with various milling times (12, 24, and 48 h) after each process. Regardless of milling time, the FCC phase of the CoCrFeMnNi HEA matrix was maintained after ball-milling without any reaction between the metallic elements of the HEA powder and the Y₂O₃ particles. The peaks for Y₂O₃ particles broadened with increasing milling times because the Y₂O₃ particles shattered.

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Figure 3

XRD patterns of the ODS composites with various milling times [(a) 12 h, (b) 24 h, and (c) 48 h] after each process.

After hot pressing, no significant phase transformation was observed for samples produced using powder milled for 12 and 24 h. However, the sigma phase and Mn oxide appeared in the sample produced using powder milled for 48 h. Longer milling time generated a greater number of dislocations, which can simulate atomic diffusion and consequently accelerate the formation of new phases.

After the heat treatment, the intensity of the peaks corresponding to the BCC phase, Cr oxide, and Mn oxide was observed in all ODS composites as in situ second phases during the process, as shown in Figure 3. The sigma phase, appeared during the hot-pressing of the powder milled for 48 h, disappeared during the heat treatment at 800°C for 12 h. The sigma phase is stable in the CoCrFeMnNi HEA at temperatures lower than 800°C and can be decomposed at temperatures over 800°C [25,26]. Therefore, the sigma phase disappeared after the heat treatment at 800°C for 12 h. On the other hand, the BCC phase and oxides were precipitated from the CoCrFeMnNi HEA matrix. Mechanically milled powders generally have a high oxygen level because of their large specific surface area [27]. This provides an oxygen source for the in situ formation of various metallic oxides during the heat treatment.

Figure 4 shows the XPS Cr2p and Mn2p spectra of the heat-treated ODS composites with powders milled for 12, 24, or 48 h. The type of oxides corresponding to the matrix elements (Co, Cr, Fe, and Ni) of the ODS composites was similar, except for Mn. Cr oxides were precipitated as CrO₂ and Cr₂O₃ within the ODS composites after the heat treatment at 800°C for 12 h, as shown in Figure 4(a)–(c) and in the EDS maps of the heat-treated ODS composites (Supplementary Figure 2). The peak corresponding to MnO was not observed in the ODS composite with powder milled for 48 h (Figure 4(d)–(f)). The complex ODS composites were fabricated with oxides that had not only ex situ Y₂O₃ but also various in situ metallic oxides.

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Figure 4

XPS Cr2p and Mn2p spectra of the ODS composites with various milling times ((a,d) 12 h, (b,e) 24 h, and (c,f) 48 h) after heat treatment.

Figure 5 shows TEM bright-field images of the ODS composites with powders milled for 12, 24, and 48 h. The average grain size of the composites was 224.91 ± 66.51, 158.23 ± 19.02, and 103.24 ± 14.72 nm, respectively. The milling energy can be expressed by E_mill = κω³t, where κ is a constant related to milling, ω is the speed of the mill rotation, and t is the milling time [28,29]. Therefore, the total energy during ball milling increases with increasing milling time [30]. Consequently, the grain size of the milled powder was refined with increasing milling time because the total energy can be stored in the form of dislocations that may act as a source of dynamic recrystallization. The ODS composite with powder milled for 48 h received significantly more impact energy from the milling media during the ball-milling process than the other composites. Therefore, its particle and grain sizes were smaller than those of other composites before the hot pressing and heat treatment. Consequently, the grain size of the ODS composite decreased as the milling time increased from 12 to 48 h after the hot-pressing and heat treatment. Significant grain growth was not observed during the hot-pressing possibly because of the presence of multiple oxides.

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Figure 5

TEM bright-field images of the ODS composites with various milling times [(a,b) 12 h, (c,d) 24 h, and (e,f) 48 h] after heat treatment.

Figure 6(a)–(c) shows the true stress–strain curves of the ODS composites for various milling times (12, 24, and 48 h). The yield strengths of the heat-treated ODS composites were ∼0.82, ∼0.92, and ∼1.40 GPa, respectively. Strengthening in these HEA-ODS alloys can be accompanied by three main mechanisms: (i) solid solution hardening, (ii) grain boundary strengthening in the HEA matrix, and (iii) dispersion hardening by multiple oxides. First, according to the lattice distortion effect of high-entropy alloys based on atomic size differences, the mechanical properties of the ODS composites improved because the HEA matrix structure of the ODS composite was maintained after every process. Therefore, the basic solid solution strengthening of HEA was sufficient in the ODS composites.

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Figure 6

True stress–strain curves of the ODS-HEA composites with various milling times [(a) 12 h, (b) 24 h, and (c) 48 h] and (d) comparison with mechanical properties of other HEAs.

The yield strength of the ODS composites increased with increasing milling time. This mainly originated from the grain size reduction during milling and the simultaneous differences in the oxide type and distribution within the ODS composites. The ODS composite using 48 h-milled powder possessed more MnO₂ than MnO, thus being tougher. In addition, the oxides in the ODS composite were distributed more homogeneously, including ex situ Y₂O₃ nanoparticles. Moreover, the grain of the ODS composites was refined as the milling time increased from 12 to 48 h. Consequently, the yield strength of the ODS composite was the largest for the sample with 48 h milling than for the other composites.

However, the compressive strength of the ODS composite with powder milled for 48 h was smaller than that for the composite with powder milled for 24 h. In situ second phases that appeared during the hot pressing or heat treatment, such as the BCC phase and oxides (i.e. MnO₂), occurred due to the brittleness of the composites. Therefore, the mechanical properties of the ODS composites were affected by the type of Mn oxide because MnO₂ (shear modulus (G) = 66.5 GPa, bulk modulus (B) = 102.5 GPa, and B/G = 1.54) were more brittle than those with MnO (G = 64.5 GPa, B = 143 GPa, and B/G = 2.22). The composite milled for 48 h presented a high fraction of in situ second phase, which was brittle because the composite acquired many nucleation sites during the ball-milling process. Therefore, the composite milled for 48 h was predicted to be more brittle than the other composites. It presented the smallest elongation, and it fractured before the strength reached ultimate compressive stress. Figure 6(d) compares the mechanical properties of the ODS composite of powders milled for 24 h in this study and other HEAs and their composites [10,18,31–37]. The results indicate that the ODS composite milled for 24 h presented a remarkable combination of strength and ductility.

Conclusions

CoCrFeMnNi-based ODS composites were fabricated using powder metallurgy with various milling times (12, 24, and 48 h) followed by heat treatment. Longer milling times led to a greater reduction of grain sizes in the HEA matrix and a more uniform dispersion of Y₂O₃ in the matrix. The ODS-HEA composites included both ex situ Y₂O₃ nanoparticles and in situ oxides containing Cr and Mn. The 48 h-milled ODS composites exhibited a high yield strength of ∼1.40 GPa, which was attributed to the dispersion strengthening by the nano-size Y₂O₃ particles and in situ oxides, grain boundary strengthening, and significant lattice distortion of the HEA matrix. However, the formation of brittle MnO₂ and the BCC phase led to a poor elongation of ∼7.76%. On the other hand, the 24 h-milled ODS composites exhibited a good balance of strength (∼920 MPa) and ductility (∼17.87%).