It is known that grain refinement is a very effective method to improve the strength and toughness of aluminum and its alloys [1,2]. The numerous refinement methods have been developed, including chemical refinement, rapid cooling, mechanical vibration, electromagnetic stirring and ultrasonic vibration [[3], [4], [5], [6]]. Among them, chemical refinement is commonly used in industrial production because of its high efficiency and economic scale [7,8]. Al-Ti-B master alloys are widely applied to refine aluminum and its alloys, but TiB₂ particles easily aggregate, and Ti reacts with impurity elements such as Zr, Cr, V, and Mn in aluminum melt, which can poison the refiners leading to a significant decrease in the refining ability of Al-Ti-B master alloys [[9], [10], [11]].

Over the course of researching the grain refinement of aluminum alloys, some new refiners have been developed. Doheim et al. [12] synthesized an Al-Ti-C master alloy in order to refine the pure aluminum and its alloys by reaction of a compacted mixture of K₂TiF₆ and graphite with molten aluminum. The result showed that the refining effect of Al-Ti-C master alloys was more significant than that of Al-Ti-B master alloy for aluminum and its alloys. Nie et al. [13] used a melt reaction method to prepare an Al-Ti-C-B master alloy with a uniform microstructure whose refinement efficiency did not significantly decrease within 60 min. In addition, it was stated that the addition concentration of Ti is less than 0.15 wt.% in aluminum and its alloys [2]. Wang et al. [14] reported that the mean size of α-Al grains was refined to 40 μm when 2.0 wt.% Al-5Ti-0.25C refiner was added to 6063 alloy. The average size of α-Al grains decreased 40 μm as the concentration of Al-5Ti-1B refiner was increased to 3.0 wt.%. However, further increasing the concentration of refiners (Al-5Ti-0.25C or Al-5Ti-1B), the size of α-Al grains tended to coarsening. Wang et al. [15] prepared an AlN-TiN/Al refiner to refine pure aluminum by a plasma jet method. The result indicated that the tensile strength, yield strength, and micro-hardness of pure aluminum were improved, except for the elongation decreased with increasing addition concentration of AlN-TiN/Al. Watanabe et al. [16] studied the refining efficiency of α-Al grains in pure aluminum by adding Al_2.5Cu_0.5Ti particles with an L1₂ structure. The prepared Al_2.5Cu_0.5Ti refiner significantly refined α-Al grains of pure aluminum, and the grain refinement effect faded more slowly as the volume fraction of the heterogeneous nuclei increased.

Moreover, previous literature reported that alloying is an effective method to refine grains through adding Ti, Nb, Ni and Cr elements [17]. Pongen et al. [18] reported that the Al-3Co master alloy as a grain refiner could refine the microstructure from coarse dendrites to fine equiaxed crystals in aluminum alloys. Sha et al. [19] showed that the addition of Co significantly changed the morphology of intermetallic compounds from small block shapes to comparatively coarse dendrites or fish-bone structure, and improved the mechanical properties of aluminum alloys. Tian et al. [20,21] studied the influence of Fe on the microstructure of aluminum-silicon alloys. The result indicated that (AlSiFe) clusters enhanced the nucleation potency and refined the α-Al and silicon phases. Yu et al. [22] showed that adding Ni-Si master alloy could decrease the density of α-Al in an aluminum-silicon alloy and refine primary silicon by reducing the undercooling. In addition, Liu et al. [23] proposed that the addition of Ni could refine the α-Al dendrites of an Al-Zn alloy containing Mg and Cu, and simultaneously improved the strength and elongation. These reports showed that the separate addition of Co, Fe, Ni, and Ti could refine microstructure and enhance the strength of the resulting alloys. However, it was unclear whether the grain refining effect of aluminum and its alloys through complex addition of these elements was much more obvious than that of separate addition of these elements. Li et al. [24] first studied the effect of trace amounts of AlCoCrFeNiTi HEA on the eutectic silicon and α-Al of hypoeutectic Al-7Si alloy. It was indicated that the addition of 0.2 wt.% HEA effectively refined the eutectic silicon structure from a coarse lamellar shape to fine rod and granularity. The ultimate tensile strength and ductility of the Al-7Si alloy were also significantly improved. Meanwhile, the size of α-Al grains obviously reduced when 0.2 wt.% HEA was added, but the dendritic morphology of α-Al was still maintained.

Based on the previous studies, it is very valuable to investigate the effect of combined addition of Co, Fe, Ni, Cr and Ti elements on the microstructural evolution and mechanical properties of aluminum alloys. Therefore, we prepared AlCoCrFeNiTi high-entropy alloy (HEA) by vacuum arc melting to realize simultaneous addition of these elements to an aluminum melt. Moreover, in order to study the influence of HEA on the morphology and size of α-Al, precipitates and mechanical properties of aluminum alloys, commercial pure aluminum was selected in the research project, and the HEA content was varied from 1.0 wt.%, 2.0 wt.%, to 3.0 wt.%. The refinement and strengthening mechanisms were also discussed in detail.

High-purity metals of Al, Co, Cr, Fe, Ni, and Ti were used as raw materials and melted together to prepare AlCoCrFeNiTi HEA through vacuum arc melting furnace under an Ar atmosphere. These prepared HEA need to be remelted at least 5 times to ensure a uniform chemical composition and microstructure.

First, commercial pure aluminum (99.8%) was heated to 800 °C in a Si-C rod melting furnace. Next, the preheated HEA were added to pure aluminum melt and held for 60 min and stirred vigorously every 5 min to insure a homogeneous chemical composition. After that, the getter hexachloroethane (C₂Cl₆) was wrapped in aluminum foil and pressed into the alloy melt to remove gases and inclusions. Finally, the temperature of melt was decreased to 720 °C, and then the melt was poured into a conical permanent steel mold (a upper diameter of 18 mm, a lower diameter of 13 mm, and a length of 140 mm) preheated up to 200 °C. Table 1 shows nominal chemical compositions of the aluminum with different HEA content of 1.0 wt.%, 2.0 wt.%, and 3.0 wt.%, respectively.

The specimens were cut from the middle position of each casting bar, and metallographic specimens were etched by 10 vol.% hydrofluoric acid to distinguish the morphology of the α-Al and other precipitates. The microstructure was analyzed using an FEG 450 field emission scanning electron microscope (SEM) equipped with an energy dispersive X-ray spectroscopy (EDS). The mapping distribution of elements was obtained using electron probe micro-analyzer (EPMA). Meanwhile, the precipitates were investigated using X-ray diffraction (XRD) (D/max 2400) and transmission electron microscopy (TEM) to analysis these phases in the as-cast samples. TEM samples were prepared using a Gatan Precision Ion Polishing System.

The effects of different AlCoCrFeNiTi HEA concentrations on the strength and ductility of industrial pure aluminum were examined by conventional tensile tests. The samples of tensile test were machined in terms of the ASTM E8/E8M-13a standard. The gauge length is 25 mm and the diameter is 5 mm. Tensile tests were performed at a rate of 1 mm/min using a Shimadzu AGS-X mechanical tester at room temperature to obtain the value of UTS, YS, and El. Moreover, the morphology of fracture was observed using SEM to analyze the fracture mode.

Fig. 1 shows the XRD pattern of HEA. The XRD analysis of HEA shows that the microstructure of alloy is comprised of two phases with body-centered cubic (BCC) structure (denoted as B2 and A2) and other two phases with complex cubic structure (denoted as D0₃ and A12).

The XRD pattern of the sample with 3.0 wt.% HEA is shown in Fig. 2. It can be seen from Fig. 2 that the sample contains Al, Al₃Ti, Al₃Ni, and some unknown phases. Fig. 3 shows the macrostructural evolution of aluminum before and after refinement through adding HEA. The α-Al exists as coarse columnar grains of pure aluminum, and completely cover the sample section from the center to the edge, as shown in Fig. 3(a). When 1.0 wt.% HEA is added to pure aluminum, a few columnar crystals present at the edge of the specimen. However, most coarse columnar crystals are refined into equiaxed grains, as shown in Fig. 3(b). As the addition level of HEA increases to 2.0 wt.% and 3.0 wt.%, the grains are further refined and the grain boundaries are nearly indistinguishable, as shown in Fig. 3(c, d). The SEM images of aluminum with different HEA contents in Fig. 4 show that the microstructure is composed of a gray α-Al and a white precipitate distributing interdendritic regions of α-Al. Fig. 4(a) shows that the microstructure of unrefined pure aluminum is coarse columnar crystals and the mean size of α-Al grains is about 374 μm. However, when 1.0 wt.% HEA is added to the molten pure aluminum, the morphology of α-Al transforms from coarse grains to relatively fine dendrites, as shown in Fig. 4(b). Increasing the addition amount of HEA to 2.0 wt.% further refines the α-Al dendrites, as shown in Fig. 4(c). When the concentration of HEA is 3.0 wt.%, the α-Al grains are significantly refined from dendrites to fine equiaxed crystals, and the average size of α-Al grains decreases to 27 μm. Furthermore, plate-like compounds with an average length of 30 μm and an average width of 2 μm distribute throughout the α-Al matrix, as indicated by the black arrows in Fig. 4(d).

In order to study the existence form of these elements in aluminum alloy, the mapping analysis of the aluminum with 3.0 wt.% HEA are obtained by EDS, as shown in Fig. 5. According to the mapping analysis of elements, it is obviously observed from Fig. 5(a) that the distribution of Al and Ti elements well matches that of the massive plate-like compounds. Fig. 5(b) displays a fine compound with a size of approximate 3 μm. Moreover, it is obvious that the distribution of Al and Ni elements closely matches to the fine particles. Combining the XRD analysis (Fig. 2) with the mapping analysis (Fig. 5) of aluminum alloy with 3.0 wt.% HEA, this shows that the plate-like compounds are Al₃Ti, and the fine particles are Al₃Ni.

Fig. 6 demonstrates the morphological evolution of precipitates distributing interdendritic regions of α-Al when 1.0 wt.%, 2.0 wt.%, and 3.0 wt.% HEA is added to pure aluminum melt, respectively. Fig. 6(b, d and f) are magnified images of Fig. 6(a, c and e), respectively. It can be seen that the dense nano-phases precipitate and distribute interdendritic regions of α-Al in the solidification after the addition of HEA to pure aluminum. Moreover, the proportion of the dense nano-phases gradually increases and the area fraction of α-Al reduces as the HEA concentration increases. In addition, it can be found that the nano-phases are parallel to the observation surface, as shown in Fig. 6(d, f). Fig. 7 shows the morphology of nano-phases located in different regions. It is observed from Fig. 7(a, b) that a certain angle is formed between the nano-phases and the observation surface of the aluminum sample with 3.0 wt.% HEA. That is to say, the growth direction of nano-phases has distinct difference during the solidification, as the schematic diagram of the nano-phases growth direction shown in Fig. 8. The statistical results of the length and diameter on the nano-phases are shown in Fig. 9. It is worth noting that the length of nano-phases gradually increases with increasing HEA content, and the average lengths are approximately 2568 nm, 4372 nm, and 6907 nm, respectively. Moreover, the nano-phases diameter gradually decreases with the increase of HEA concentration, with mean diameters of 112 nm, 103 nm and 92 nm, respectively.

Fig. 10 shows TEM dark field images and corresponding selected area electron diffraction (SAED) patterns, which show the morphology and structure of the nano-phases in aluminum containing 3.0 wt.% HEA. The result further demonstrates that some nano-phases is parallel to the observation surface, as shown in Fig. 10(a). However, there is a certain angle formation between other nano-phases and the observation surface, as shown in Fig. 10(b). In addition, the SAED pattern shows that the nano-phase belong to body-centered cubic structure. Fig. 11 presents the mapping analysis of various elements of nano-phases. It is obviously observed that the nano-phases consist of Al, Fe, Co, and Ni elements. It is therefore inferred that the dense nano-phases are AlFeCoNi compounds with body-centered cubic structure. In addition, it is noticeable that there is no Cr-rich intermetallic compounds and Cr is uniformly distributed in the α-Al matrix, as shown in Fig. 11.

The commonly accepted refinement mechanism of α-Al grains in aluminum and its alloys involves the effective heterogeneous nucleation and solutes segregation within the molten alloy [25,26]. Many fine particles forming through reaction act as heterogeneous nucleation sites and some solutes give rise to segregation in the front of solid-liquid interface refining the α-Al grains. To study the refinement mechanism of α-Al grains due to the addition of HEA, it is necessary to analyze the existence and distribution of elements. Fig. 12 shows the EPMA element distribution maps of Al, Ti, Cr, Co, Fe and Ni in the alloy containing 3.0 wt.% HEA. The concentration of Ti is 0.48 wt.% (Table 1) when 3.0 wt.% HEA completely dissolve in aluminum melt. The back-scattered electron image (Fig. 12(a)) shows that a white particle exists at the center of the α-Al grain with a size of approximate 2 μm. Observing the mapping analysis, it is found that the white particle is composed of Al and Ti, as shown in Fig. 12(b, c). This indicates that stable Al₃Ti particles are preferentially formed prior to the α-Al and act as heterogeneous nucleation sites of α-Al grains in the solidification when the Ti-containing HEA is added to pure aluminum. The refining effect of these refiners including Al₃Ti (a =0.3846 nm, c =0.8594 nm), TiC (a =0.4328 nm), TiB₂ (a =0.3038 nm, c =0.3239 nm), and AlB₂ (a =0.3009 nm, c =0.3262 nm) on α-Al (a =0.4049 nm) is predicted by an edge-to-edge matching model [27]. The numerical results indicated that the grain refining efficiency of Al₃Ti is much better than that of TiC, TiB₂, or AlB₂ because the misfit between Al and Al₃Ti is smallest in comparison with other refiners. The previous investigations also indicated that the refining effect of Al₃Ti is better than that of TiB₂ for aluminum and its alloys [[28], [29], [30], [31]]. HEA containing Ti dissolves when it is added to pure aluminum melt in our experimental condition, and Ti would form fine Al₃Ti particles reacting with Al in the aluminum melt. Based on the analysis above, Al₃Ti particles are shown to act as the heterogeneous nucleation substrates for α-Al grains.

Moreover, it is worth noticing that the Ti-rich regions are formed within the α-Al grains and surround the Al₃Ti particles, as shown in Fig. 12(c). The formation of a Ti-rich monolayer surrounding TiN significantly improves the refining efficiency of the α-Al grains when TiN/Ti is added into pure aluminum [32]. The formation of Al₃Ti layer located at the Al/TiB₂ interface in the Al/Al-Ti-B system can improve the nucleation potency of α-Al grains, and segregated Ti atoms in the aluminum generate constitutional undercooling and activate more substrates to participate in nucleation [26,28,33]. The formation of a Ti-rich zone at the Al/TiC interface also improves the nucleation efficiency of TiC for α-Al grains [34,35]. In addition, Easton and StJohn [36,37] investigated the relationship between grain size (d) and solute content (defined by the growth restriction factor, Q), the potency and number density of nucleation particles. The grain size (d) is expressed as follows:

where c0 is the volumetric number density of nucleant particles; f is the fraction of these particles that are activated; D is the diffusion coefficient; v is the growth velocity; ΔTn is the undercooling for active nucleation, and Q is the growth restriction factor (GRF) [38]:

where c0 is the initial solute concentration in the liquid,ml is the liquidus slope, and k is the solute partition coefficient. Eq. (1) shows the reciprocal relationship between the grain size and the growth restriction factor (Q). While the growth restriction factor can evaluate the grain refining efficiency of solutes. That is to say, the grain refining efficiency is related to the growth restriction factor (Q). It can be seen from Eqs. (1) and (2), the grain size is inversely proportional to growth restriction factor (Q). The Q value is the largest for Ti among these alloying elements such as V, Zr, Nb, Cr and Mg in aluminum alloys [39]. Ti acts as a powerful segregating solute and generates constitutional undercooling in the front of growing crystals, which can activate more nucleation particles leading to grain refinement. Consequently, the α-Al grains are refined through heterogeneous nucleation and solute segregation.

Cr does not form intermetallic compounds (e.g., Al₇Cr) and is uniformly distributed in the aluminum matrix when the concentration of Cr is less than 0.6 wt.%, as shown in Fig. 11, Fig. 12(d). This phenomenon is consistent with the microstructure of rapidly solidified Al-Cr alloy [40]. Herein, the Al-Cr alloy is self-inoculating binary system, and the addition of Cr can effectively refine α-Al grains of aluminum and its alloys. Moreover, the average grain size quickly decreases as the Cr content increases [41,42]. According to the analysis, Cr can form the supersaturated solid solution. This shows that the grain refinement cannot be explained by the growth restriction factor (GRF) because the growth restriction effect is associated with the solute segregation in the front of solid-liquid interface. In addition, it is known that the Cr is peritectic forming element in aluminum, and the maximum solubility of Cr in aluminum matrix is 0.29 wt.% at 661.5 °C. When the concentration of Cr exceeds the max solubility, peritectic particles (Al₇Cr) are formed prior to α-Al and peritectic reaction occurs during the solidification.

In addition, the mapping analysis in Fig. 12 shows that the Co, Fe and Ni has limited solubility in the α-Al matrix except Ti and Cr elements, causing them to be pushed to the front of solid-liquid interface of α-Al growth. A solute-rich layer is formed, which leads to the formation of constitutional undercooling in the front of solid-liquid interface during the solidification. According to the analysis of microstructure, the area fraction of nano-phases increases with increasing the concentration of Co, Fe and Ni elements. The result shows that lots of solute atoms aggregate in the front of solid-liquid interface. Chen et al. [43] reported that the atomic aggregation can increase constitutional undercooling and restrict the crystal growth ahead of solid-liquid interface in the solidification. Moreover, solute segregation can quickly form constitutional undercooling and more particles are activated for heterogeneous nucleation substrates. Hence, the solute segregation of Co, Fe and Ni in the front of solid-liquid is favorable to formation of constitutional undercooling and activating more nucleation particles.

According to the analysis above, the α-Al grains refinement is mainly attributed to three factors: (1) Al₃Ti particles act as heterogeneous nuclei of α-Al grains; (2) Cr is a self-inoculating peritectic solute in Al-Cr binary alloy; (3) The segregation of Co, Fe and Ni generating constitutional undercooling in the front of solid-liquid interface.

Fig. 13 displays the UTS, YS and El of pure aluminum with different HEA concentrations. Compared with unrefined pure aluminum, the UTS and YS are significantly improved, but the El gradually is decreased as the HEA content increases. The UTS is improved by 145% from 62 MPa to 152 MPa, YS is increased by 173.8% from 42 MPa to 115 MPa and the El is decreased by 33.3% from 39% to 26% when 3.0 wt.% HEA is added.

Fig. 14 demonstrates the tensile fracture morphology of aluminum with the different HEA concentrations. Fig. 14(a) shows the fracture surface of pure aluminum, which is composed of many deep dimples, corresponding to the excellent elongation of pure aluminum. However, as the HEA concentration increases, the dimple depth becomes shallow and the size decreases, as shown in Fig. 14(b-d). This is due to increasing the number of stiff and brittle Al₃Ti and dense nano-phases. Hence, the dimples on the fracture surface become shallower and smaller, and the El significantly reduces.

As illustrated in Fig. 13, the mechanical properties of aluminum are related to the amount of added HEA. Moreover, the UTS and YS are gradually improved with increasing HEA concentration. As we known, the strength and toughness of aluminum and its alloys depend on the microstructure. Based on the refinement mechanism of α-Al grains by HEA, it can be deduced that the strengthening of aluminum via HEA addition results from solid solution strengthening, dispersion strengthening and nano-phases strengthening. According to the analysis of Fig. 11, Fig. 12, Cr atoms absolutely dissolve into the Al solid solution, which causes severe lattice distortion. The atoms segregate in the vicinity of dislocation to lower the energy of lattice distortion. The atoms pining may increase resistance of dislocation movement. Moreover, the strengthening is related to the solute content in the solid solution, which is described as follows [44]:

where Δτss is the change in yield stress due to solid solution strengthening, A is the constant and C0 is the concentration of the solute (wt.%). The Cr concentration increases from 0.17 wt.% to 0.52 wt.% as the increasing HEA content (Table 1).

Fig. 4(b) shows that fine Al₃Ni particles are distributed in the aluminum matrix. It was reported that Al₃Ni is responsible for strengthening due to its high Young’s modulus (116-152 GPa), hardness (HV_0.01 = 5130 MPa) and tensile strength (2160 MPa) [23]. Therefore, the Al₃Ni dispersoid can serve as pinning particles that hinder dislocation glide and the strength is improved in tensile. Fig. 6 shows that the area fraction of the nano-phases gradually increases, and the fraction of Al solid solution decreases with increasing HEA content. The dense nano-phases can make the distribution of deformation and stress concentration more uniform, which may suffer much more deformation [45]. In addition, the strength of alloys can be significantly enhanced when the lamellar spacing of eutectic alloys is decreased [46]. Accordingly, the UTS and YS are improved when the addition level of HEA is 1.0 wt.%, 2.0 wt.% and 3.0 wt.%, respectively.

Additionally, He et al. [47] reported that the Hall-Petch coefficient (k) is 0.04 MPa m^1/2 in aluminum alloys. The increment in the calculated YS is only 5.6 MPa when the grain size of α-Al decreases from 374 μm to 27 μm. Therefore, fine grain strengthening can be ignored in the investigation. It is noted from the Fig. 13 that the El gradually decreases when the addition level of HEA is increased from 1.0 wt.% to 3.0 wt.%. According to the analysis of phases, the microstructure is composed of plate-like Al₃Ti, fine Al₃Ni particles, dense nano-phases and α-Al solid solution when HEA is added into aluminum melt. The solid solution strengthening, dispersion strengthening and nano-phases strengthening are due to pinning effect that inhibits the dislocation movement in tensile deformation. Hence, the UTS and YS are improved, but the El is gradually decreased as the HEA concentration increases.

The effects of AlCoCrFeNiTi HEA on the microstructure and mechanical properties of pure aluminum were studied, and the major conclusions can be drawn as follows:

(1)The addition of HEA can effectively refine the α-Al grains into fine equiaxed crystals. Moreover, a higher HEA concentration results in a finer α-Al grains. AlCoCrFeNiTi alloy is an effective grain refiner for aluminum and it alloys.

(2)The addition of HEA results in precipitates of Al₃Ti, Al₃Ni and dense nano-sized AlFeCoNi phases. The length of nano-phases increases from 2568 nm to 6907 nm, but the diameter decreases from 112 nm to 92 nm.

(3)Compared with the pure aluminum, UTS and YS are improved, whereas El is decreased with increasing the HEA concentration. The UTS is enhanced by 145.2% from 62 MPa to 152 MPa, YS is increased by 173.8% from 42 MPa to 115 MPa, but the El is decreased by 33.3% from 39% to 26% when the addition of HEA increased to 3.0 wt.%.

This work was financially supported by the National Natural Science Foundation of China (No. 51561021), the State Key Laboratory of Advanced Processing and Recycling of Nonferrous Metals, Lanzhou University of Technology (No. SKLAB02019007), and the National Innovation Training Program of College Students of China (No. DC2019165 DC2019161).