High-entropy alloys (HEAs) [1,2,3,4,5,6,7], as an emerging class of metallic alloys, exhibit great potentials in structural applications because of their excellent mechanical properties. HEAs are usually defined as consisting of five or more principal elements in equiatomic or near equiatomic ratios [8, 9]. Based on this design concept, HEAs often possess high configurational entropies, facilitating the formation of a single solid solution phase in a moderate range of enthalpy of mixing or atomic size differences [10]. Recent studies have demonstrated that unique structural features of HEAs can provide an excellent combination of high strength, high plasticity, high hardness and outstanding wear resistance, especially at cryogenic temperature, which enable HEAs promising for wide cryogenic applications in the fields like liquid gas storage, space exploration and superconductivity [11,12,13,14]. For instance, Otto et al. [11] reported a simultaneous improvement in both strength and ductility of CoCrFeMnNi HEA while reducing the tensile temperature, and attributed such phenomenon to the formation of deformation nanotwins upon tensile loading. Fu et al. [12] clarified the effect of testing temperature on the microstructural evolution and tensile properties of CoCrFeMnNi HEA via in situ synchrotron-based high-energy X-ray diffraction tensile tests from 298 down to 123 K and revealed the strengthening mechanisms in different temperature conditions. Accordingly, the deformation mechanism of HEAs transferring from dislocation slipping to deformation twinning is of essence to gain an outstanding mechanical performance at cryogenic temperature.

To date, the fabrication of HEAs mainly depends on conventional casting methods, such as vacuum arc-melting and drop casting, but the dimensions of the bulk ingots are limited to the processing method and cooling rate [15, 16]. Moreover, serious compositional segregation in the ingots commonly exists, and thus, the ingots are generally subjected to subsequent heat treatment to homogenize the microstructure and constituent [17]. To address this issue, laser melting deposition (LMD), as a low-cost and flexible additive manufacturing (AM) technology, has been applied to the manufacturing of HEAs [18,19,20]. LMD is a process that builds metallic components with a shape satisfying the requirements for geometry complexity personal customization, in a form of point by point and layer by layer directly from computer-aided design (CAD) models without specific tooling and manual intervention [21, 22]. Ultrafast cooling rate (1 × 103–1 × 104 K·s−1) of LMD process can endow the parts with extraordinary non-equilibrium microstructures and avoid elemental segregation, thus obtaining a high performance HEA component [23]. Gao and Lu [24] successfully manufactured bulk CoCrFeMnNi HEA with an excellent combination of high ultimate tensile stress and high elongation via laser 3D printing technology. Chew et al. [25] studied the microstructural characteristic of columnar to equiaxed transition, and remarkable mechanical properties of high yield strength and ultimate tensile strength of CoCrFeMnNi HEA fabricated by laser aided additive manufacturing technique, and quantitatively analyzed the strengthening effect of refined grains. Although more attentions have been paid on the tensile properties of AM-fabricated HEA components at room temperature, their mechanical properties under cryogenic conditions have been rarely investigated, especially for wear resistance of HEAs.

Inspired by earlier findings, the main objective of the present work is to fabricate bulk CoCrFeMnNi HEA via LMD technique and investigate the tensile properties and wear resistance of the as-built HEA samples at room temperature and cryogenic temperature. It is expected that the obtained results can offer more insights for better understanding the mechanical properties of additively manufactured CoCrFeMnNi HEA at cryogenic temperature and thus widening its industrial applications in extreme conditions.

Rectangular samples with a dimension of 30 mm × 30 mm × 20 mm were fabricated via LMD on an LDM-8060 workstation with a YSL-6000 fiber laser. Pre-alloyed CoCrFeMnNi HEA powders with an average particle size of ∼75 μm were used for the LMD process as the raw material, which exhibited spherical or near spherical shape with a few satellite particles, and homogeneous elemental distribution, as shown in Fig. 1a. The chemical compositions of HEA powders are 20.11% Co, 19.32% Cr, 20.45% Fe, 20.33% Mn and 19.79% Ni. The HEA powders were continuously injected into the molten pool by argon flow at a feeding rate of 10 g·min−1. The 45# medium carbon steel plate was selected as the base substrate due to its comparable expansion coefficient with the studied CoCrFeMnNi HEA. The steel plate was carefully ground and ultrasonically cleaned in ethyl alcohol and deionized water sequentially for 10 min in each step prior to the LMD process. The HEA samples were fabricated by a set of processing parameters including a laser beam diameter of 3 mm, a laser power of 1400 W, a scanning rate of 600 mm·min−1 and a layer thickness of 0.5 mm, and in an inert argon atmosphere with the oxygen concentration below 10 × 10–6. An overlapping fraction of 30% between the adjacent laser tracks was set to obtain a homogeneous and continuous deposition.

Fig. 1
figure 1

a SEM image of CoCrFeMnNi HEA powders and corresponding EDS analysis showing elemental distribution of yellow box; b XRD patterns of CoCrFeMnNi powders and as-built HEA samples; c SEM image of as-built HEA samples and corresponding EDS analysis showing elemental distribution of yellow box

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The phase structure was identified via X-ray diffraction (XRD, Empyrean) with a CuKα radiation source. The microstructures were characterized by using a scanning electron microscope (SEM, Zeiss-SUPRA55) equipped with an energy-dispersive spectrometer (EDS), and a transmission electron microscope (TEM, FEI Talos F200x). Mechanical properties at room temperature and cryogenic temperature were characterized in terms of microhardness, wear and tensile tests. Microhardness tests were performed using a Vickers indentation tester (HVS-1000A) at a load of 1.96 N and a dwell time of 15 s. Wear tests were carried out using a pin-on-disc tribometer (YTN-TB) at a normal load of 10 N, a sliding speed of 0.8 m·s−1 and a sliding time of 40 min. The worn surfaces were examined using SEM to study the wear mechanism. Dog bone-shaped tensile samples with a gauge length of 7 mm, a width of 1 mm and a thickness of 1 mm were electric discharge machined from the as-built HEA samples, and carefully grounded and polished to a mirror surface by SiC paper. All of the uniaxial tensile tests were performed on a universal testing machine (AGXplus) at a nominal strain rate of 1 × 10–3 s−1. The microstructure of the HEA samples after tensile tests was examined via TEM, and TEM foils were cut from the uniformly deformed regions of the HEA samples after tensile deformation, carefully grounded down to ~ 50 μm thick and then thinned by twin-jet electro-polishing to ensure the electron transparent. All of the mechanical samples were cooled by liquid nitrogen to desired cryogenic temperature and stabilized for 10 min before the mechanical tests. All the mechanical tests were repeated at least five times to ensure data reproducibility.

Figure 1b shows XRD patterns of CoCrFeMnNi powders and as-built HEA samples. As seen, both the as-built HEA samples and the raw powders possess a typical single-phase fcc structure, suggesting that no new phases were formed and fcc solid solution structure remains unchanged during the LMD process.

SEM image and EDS analysis of the as-built HEA samples are shown in Fig. 1c. Few pores or cracks can be observed in the as-built HEA samples. During the LMD process, a local high temperature caused by heat input with high volumetric density can maintain for a relatively long duration, thus providing sufficient time to enable the gas escape from the molten pool [26]. In addition, the microstructure of the as-built HEA samples consists mainly of dominant columnar grains and a small number of equiaxed grains. The columnar grains grow along the deposition direction, which is usually parallel to the heat flow direction [27]. Also, EDS analysis indicates that the principal constituent elements of Co, Cr, Fe, Mn and Ni are homogeneously distributed in the as-built HEA samples, suggesting that no elemental segregation occurs during the LMD process.

To examine the tensile properties of the as-built HEA samples at room temperature and cryogenic temperature, uniaxial tensile tests were performed at 298 and 77 K, and the representative engineering stress–strain curves are plotted in Fig. 2a. The as-built HEA samples exhibit the yield strength (σy) of (346 ± 17) and (470 ± 23) MPa, ultimate tensile strength (σUTS) of (462 ± 19) and (635 ± 31) MPa and elongation to failure (εf) of 21.1% ± 2.0% and 30.5% ± 1.7% at room temperature and cryogenic temperature, respectively. Apparently, the cryogenic temperature leads to a remarkable enhancement in both strength and plasticity of the as-built HEA samples. The fractured surfaces of the as-built HEA samples after tensile testing at 298 and 77 K are shown in Fig. 2b, c. As seen, a large number of dimples are homogeneously distributed on the fractured surfaces, which is a typical feature of ductile fracture. The fractured surface of samples tested at 77 K appears deeper and larger dimples than that at 298 K, which is beneficial to the pronounced ductility.

Fig. 2
figure 2

a Representative engineering stress–strain curves of tensile tests at 298 and 77 K and (inset) dimension of tensile test sample (unit: mm); SEM images of fractured surfaces of as-built HEA samples after tensile testing at b 298 and c 77 K; bright-field TEM images and (insets) corresponding SAED patterns along [110] zone axis of as-built HEA samples after tensile deformation at d 298 K and e 77 K

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To further reveal the deformation mechanism of the as-built HEA samples at 298 and 77 K, the microstructure of the deformed samples after tensile testing was further characterized by TEM. Figure 2d, e represents the bright-field (BF) TEM images and selected area electron diffraction (SAED) patterns of the fractured HEA samples at 298 and 77 K. It is worth noting that high-density dislocations are found in Fig. 2d, suggesting that the deformation behavior is mainly dominated by dislocation slipping at room temperature. Corresponding SAED pattern is shown in inset in Fig. 2d. Besides a high density of dislocations, the nano-sized striped deformation nanotwins are obviously observed within the matrix in the form of bundles (Fig. 2e), and both of the twin structure and orientation relationship are confirmed by the corresponding SAED pattern (inset in Fig. 2e). This implies that the twinning effect also plays an important role in the deformation behaviors of the as-built HEA samples at 77 K. It is well known that twins are expected to have a remarkable strengthening effect because of the interactions between twins and dislocations. The formation of deformation nanotwins during tensile deformation at cryogenic temperature introduces additional twin boundaries, decreases the mean free path of dislocations and thereby enhances strain hardening [28, 29]. This process is often considered as the dynamic Hall–Petch effect [30]. Besides, twin boundaries also act as sites for dislocation nucleation and accumulation [31]. Therefore, the twin boundaries formed during plastic deformation can effectively provide a source of steady strain hardening to delay the onset of necking and favor the excellent uniform elongation before fracture, thus enhancing both strength and ductility of the as-built CoCrFeMnNi HEA at cryogenic conditions [14].

Vickers microhardness tests were carried out in the as-built HEA samples at 298 and 77 K, and the corresponding microhardness values and indentation morphologies are shown in Fig. 3a, b. Notably, the microhardness value HV (214 ± 16) of the as-built HEA samples at 77 K is much higher than that HV (168 ± 11) at 298 K. As discussed above, massive dislocations will be induced into the surface of the as-built HEA samples during the hardness testing at room temperature, and profuse dislocations and newly formed deformation nanotwins will appear in the surface of samples after cryogenic temperature tests. High density of dislocations is likely to be formed by tangling and piling up near the grain boundaries, and additional twin boundaries can block dislocation movement, thereby significantly improving the microhardness of the as-built HEA samples at cryogenic conditions [32].

Fig. 3
figure 3

Microhardness and corresponding indentation of as-built HEA samples at a 298 K and b 77 K; c coefficient of friction; and d mass loss of as-built HEA samples at 298 K and 77 K

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Figure 3c shows the coefficient of friction (COF) of the as-built HEA samples. For both 77 and 298 K, COF exhibits the stable fluctuation within a certain range. The COF value of the as-built samples tested at 77 K is relatively lower than that at 298 K, suggesting an enhanced wear resistance at cryogenic conditions. The mass loss of the as-built HEA samples tested at 298 and 77 K is given in Fig. 3d. It can be seen that the mass loss of the as-built HEA samples increases with sliding time at both room temperature and cryogenic temperature. At cryogenic temperature, mass loss almost linearly increases during the wear tests, whereas mass loss dramatically increases firstly and then slightly increases at room temperature. It should be also noticed that the mass loss of the as-built HEA samples tested at 77 K is lower than that at 298 K, with the decrement from 9.67 to 3.46 mg, demonstrating the better wear resistance at cryogenic conditions.

In order to further reveal the wear behaviors of the as-built HEA samples at room temperature and cryogenic temperature, worn surfaces were examined by SEM and EDS, as shown in Fig. 4. After sliding wear at 298 and 77 K, obvious plowing grooves can be observed along the sliding direction, as shown in Fig. 4a1, b1. However, different from the plowing grooves with homogeneous width and depth on the worn surface of samples tested at 77 K, a large amount of oxide particles and delamination are distributed on the worn surface of samples tested at 298 K (Fig. 4a2). The formation of oxide particles during the wear tests at room temperature is also confirmed by EDS results, as shown in Fig. 4a3. Therefore, the wear behavior of the as-built HEA samples is mainly dominated by oxidative wear and a degree of delamination wear at room temperature. For the sliding wear at cryogenic temperature, homogeneous plowing grooves and few small delamination exist in the worn surface, and no oxide particles form (Fig. 4b2), which coincides with EDS results in Fig. 4b3.

Fig. 4
figure 4

SEM images of worn surface of as-built HEA samples tested at a1 298 K and b1 77 K; a2, b2 corresponding enlarged SEM images of red boxes in a1, b1; a3, b3 EDS analysis showing line scanning elemental results along red line in a2, b2

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Based on the above experimental observations, wear resistance of the as-built CoCrFeMnNi HEA is significantly enhanced at cryogenic temperature. On the one hand, during the wear testing under cryogenic conditions, the low oxygen concentration in the liquid nitrogen hinders the formation of oxide particles. Besides, the oxide particles on the worn surface are rarely formed in the cryogenic conditions, which are often formed due to the high flash temperature induced by dry sliding friction [33]. On the other hand, a high density of dislocations and deformation nanotwins, which are formed upon tensile loading at cryogenic temperature (TEM images in Fig. 2d, e), can significantly improve the deformation resistance of the as-built HEA samples [34]. As the tensile (Fig. 2a–c) and microhardness (Fig. 3a, b) results indicate, the ultimate tensile strength and microhardness exhibit a great enhancement at cryogenic temperature (635 MPa, HV 214) in comparison with those at room temperature (462 MPa, HV 168). In such a situation, delamination is considered to occur in the worn surface only when the plastic deformation is accumulated to reach the deformation limit of the as-built CoCrFeMnNi HEA during the sliding wear at cryogenic temperature. Accordingly, the wear resistance of the as-built CoCrFeMnNi HEA can be significantly enhanced at cryogenic temperature.

In summary, bulk CoCrFeMnNi HEA components were fabricated by LMD technique. The tensile properties, microhardness and wear resistance of LMD-fabricated HEA were investigated at room temperature and cryogenic temperature. The as-built HEA samples possess a single fcc phase and highly dense microstructures consisting of columnar and equiaxed grains, without elemental segregation. Both strength (635 MPa) and tensile plasticity (30.5%) of the as-built HEA samples exhibit a remarkable enhancement at cryogenic temperature than those at room temperature. In this condition, deformation nanotwins formed upon tensile loading at cryogenic temperature can introduce additional twin boundaries, decrease the mean free path of dislocations and thereby enhance strain hardening. Cryogenic conditions favor the enhancement in the microhardness (HV 214) of the as-built HEA samples due to the formation of dislocations and deformation nanotwins during cryogenic deformation, which can significantly improve the deformation resistance. Accompanied by improved strength and hardness, the wear resistance of the as-built CoCrFeMnNi HEA is significantly improved in cryogenic condition. The sliding wear in the liquid nitrogen can decrease the oxygen concentration and hinder the formation of the oxide particles.