Precipitation behavior of yttrium-rich nano-phases in AlCoCrFeNi_2.1Y_x high-entropy alloy

SEM images of the AlCoCrFeNi_2.1Y_{x (x = 0, 0.1%, 0.5% and 1.0%)} alloys with different Y content were shown in figures 1(a)–(d), respectively. For Y = 0, the morphology of the alloy was composed of dendrite and inter-dendrite. The dendrite was coarse, while necking was observed in the secondary dendrite. The dendrite was (Fe-Co-Cr)-rich primary phase and the liquid metal surrounded the primary dendrite nucleus, forming a (Al-Ni)-rich phase of the eutectic structure (figure 1(a)). It is noting that a fine structure with a flower-like helical morphology was formed when Y was 0.1%. With the increase of Y content, (Fe-Co-Cr)-rich phase and (Al-Ni)-rich phase of the eutectic structure grew up dependently, as shown in figure 1(b). For Y = 0.5%, 1.0% (figures 1(c) and (d), respectively), the morphology changed from a eutectic to a banded structure, and then to a dendrite structure. Meanwhile, the dendrite became coarse and the inter-dendrite became nodular, and the content of the (Al-Ni)-rich phase decreased. The morphology is related to the temperature gradient and growth rate. Dendrite structures formed when the temperature of the dendrite tip was higher than that of the eutectic alloy. The dendrite eutectic interface of (Al-Ni)-rich phase and (Fe-Co-Cr)-rich phase was in the low super cooling state of the eutectic couple zone. When the Y content was 1.0%, it resulted in a necking phenomenon at the second dendrite root as shown in figure 1(d), which indicated that Y enriched the solidification front of the eutectic structure and inhibited the growth of the inter-dendrite phase. Table 1 showed the EDS results of the AlCoCrFeNi_2.1Y_1.0% alloy. It could be seen that the microstructure of AlCoCrFeNi_2.1Y_1.0% was eutectic structure composed of (Al-Ni)-rich phase and (Fe-Co-Cr)-rich phase. A new type of (Al-Ni-Y)-rich phase (marked as C) formed with the increase of Y content. The content of Y in the (Al-Ni)-rich (marked with B) phase was higher than that in the (Fe-Co-Cr)-rich (marked with A) phase.

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Figure 1(e) displayed the XRD patterns of the AlCoCrFeNi_2.1Y_{x (x = 0, 0.1%, 0.5%, and 1.0%)} alloys. The alloys were mainly composed of BCC (B2) and FCC (L1₂) solid solutions. In all samples, Bragg peaks of FCC (FeCoNi-M) and BCC (AlNi-M) phase could be observed. The intensity of the FCC (L1₂) peaks in the (111) plane decreased with the increase of Y content, which suggested a change in the volume fraction or in the crystal orientation in FCC (L1₂). When Y ≥ 0.5%, the diffraction peaks at lower angle regions became weaker, probably indicating that Y atoms in the lattice skeleton caused orientation effect. When the Y content was 1.0%, the diffraction peak of the FCC phase in the (111) plane was weak and the peak in the (200) plane was strong.

The AlCoCrFeNi_2.1Y_{x (x = 0, 0.1%, 0.5%, and 1.0%)} samples were measured with DSC at a heating rate of 30 K min⁻¹ and a maximum heating temperature of 1350 °C, as shown in figure 1(f). When Y ≥ 0.1%, an endothermic peak occurred near 1210 °C, indicating the formation of a new phase. The curves showed an obvious endothermic peak near 1210 °C indicating the increase of the new phase when Y ≥ 0.5%. The mixing enthalpies Y-Al, and Y-Ni were calculated based on the mixing enthalpy equation (equation (1))[26], and were determined to be −38 kJ mol⁻¹ and −31 kJ mol⁻¹, respectively. A typical compound was formed by two or more elements of Y, Al and Ni for a high negative mixing enthalpy.

$$\begin{eqnarray}&&{\rm{\Delta }}{H}_{mix}\,=\,\displaystyle \sum _{i=1,i\ne j}^{n}4{\rm{\Delta }}{H}_{ij}^{mix}{c}_{i}{c}_{j}\end{eqnarray} \tag{ 1 }$$

Figure 2 showed the TEM images and selected area electron diffraction (SAED) patterns of the as-cast AlCoCrFeNi_2.1Y_{x (x = 0, 0.1%, 0.5%, and 1.0%)} HEAs. The microstructure was composed of FCC (L1₂) (Fe-Co-Cr)-rich phase (grey region), BCC (B2) (Ni-Al)-rich phase (black dendrites), and precipitates with several nanometers in size appeared in the FCC (L1₂) phase (figure 2(a)), which was similar with the results reported by Lu et al [27]. Figure 2(b) showed a TEM image of AlCoCrFeNi_2.1Y_0.5%. As could be seen in figure 2(b), a number of Y-rich nano phase appeared in the FCC (L1₂) phase marked by black arrows. The SAED patterns of the FCC phase (marked as A) and the BCC phase (marked as B) along [011] zone axis were presented in figures 2(a) and (b), respectively. The EDS results of the AlCoCrFeNi_2.1Y_{x (x = 0, 0.5%)} alloys for the FCC and the BCC phases were shown in table 2, which indicated that the microstructure of AlCoCrFeNi_2.1Y_0.5 was a eutectic structure composed of (Al-Ni)-rich B2 phase and (Fe-Co-Cr)-rich FCC (L1₂) phase. The solubility of Y into the BCC phase was higher than that of the FCC phase, and the Y content in the BCC phase was 1.0% in AlCoCrFeNi_2.1Y_0.5%.

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Figure 3 illustrated the TEM images and SAED of AlCoCrFeNi_2.1Y_1.0% before and after compressive. The growth of B2 phase was changed and the lamellar structure changed to be bamboo-like structures. Both the amount and size of nanoparticles increased when Y content increased to 1.0%. The HRTEM of the bamboo-like structure of the B2 phase (marked by the bright yellow circle) was analyzed in figure 3(a). The new phase was a BCC structure (zone axis: [0 1 2]) and formed due to super saturation precipitation and segregation of Y at the solidification front of the B2 phase. A lot of dislocations piled up at grain boundary as well as within grains. The deformation ability of the new BCC phase in bamboo joint was different from that of the B2 phase as shown in figure 3(b). The lamellar B2 phase separated from the bamboo-like BCC phase due to the segregation of Y, which might be the key factor for the decrease of mechanical properties of the alloy with the further increase of Y content.

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Figure 4 illustrated the results of compressive test of AlCoCrFeNi_2.1Y_{x (x = 0, 0.3%, 0.5%, and 1.0%)}. The compression strength and strain first increased and then decreased with increasing of Y content as shown in figure 4(a). The yield strength increased with the increase of Y content under the same strain. Double yielding behavior could be observed when Y = 0.5% and Y = 1.0% (figure 4(b)). The first and second yielding of Y_0.5% were marked and analyzed in figure 4(c). The true strain (Ln (1 + ε)) and the true stress (σ (1 + ε)) were calculated based on engineering strain (ε) and engineering stress (σ). The working hardening rate of Y_0.5% was the derivative of the true stress. The curve was shown in figure 4(d). It included three stages: (1) Stage I: the working hardening rate was increased with a high speed after elastically deformed. Beyond the elastic limit of the alloy (2.2 ± 0.2%), the first yielding occurred at 600 ± 20 MPa (marked in figure 4(c)). A slight increase in the elastic limit of the alloy could be observed due to the presence of the Y element; (2) Stage II: after the first yielding, the slope of the stress–strain curve in the strain range between 2.0 ± 0.2 and 5.0 ± 0.3%, deviated markedly from that of Stage I. The working hardening rate was grown at low speed. The second yielding occurred at 850 ± 20 MPa (marked in figure 4(c)), which caused an increase in yielding strength by 40%. (3) Stage III, the growth of working hardening rate was small and the curve was instability gradually.

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Figure 5 and 6 show the TEM images and SAED patterns of AlCoCrFeNi_2.1Y_0.5% after compression deformation. After the first yielding, the microstructure was composed of FCC (L1₂) (marked as A) and BCC (B2) phase (marked as B) as shown in figure 5. Particles with a few nanometers in size, and a large number of nano-precipitates appeared in the BCC phase. A super-lattice was formed in the SAED pattern of the B2 phase. The nano-particles were (Al-Ni-Y)-rich phase and Single Cubic (SC) structure incoherent with the B2 phase as shown in figure 5(b). The SAED pattern of the nano particle along [1 2 4] zone axis were presented in figure 5(c) based on the fast fourier transformation (FFT) result. While the nano-precipitates were Y-rich phase and FCC structure coherent/semicoherent with the FCC (L1₂) phase. The SAED pattern of the nano-precipitate along [1 1 1] zone axis was shown in figure 5(e). For Y ≥ 0.1%, the nano-precipitates formed which agreed with the DSC results as shown in figure 1(d). B2 is (Al-Ni)-riched phase. The mixing enthalpy of Al-Ni, Y-Al, and Y-Ni was −22 kJ mol⁻¹, −38 kJ mol⁻¹ and −31 kJ mol⁻¹, respectively. The HRTEM images in figure 5 and the EDS results in figure 1(d) demonstrated dispersed precipitations caused by Y. First, the solute effect [28] of Y existed in the form of atom could cause the constituent super cooling at the interface of solid-liquid, which promoted the formation of Y-rich nano-precipitates. Second, the high negative mixing enthalpies of Y-Al and Y-Ni indicated that two or more elements of Y, Al, and Ni could form compounds with low interfacial energy, and Y promoted the formation of dispersed nano-phases under the deformation stress. Dislocations (marked with black arrows in figure 5(a)) distributed near the nano-precipitates. It indicated that there was interaction between nano-precipitates and dislocations.

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The TEM image of the specimen after the second yielding, the nano-precipitates and dislocations were presented in figure 6. The nano-precipitates were incoherent with the B2 phase and dislocations distributed near the nano phase. The number of nano-precipitates and dislocations (marked with black arrows) density increased for the further deformation as shown in figure 6(a). It indicated that Yttrium further promoted the aggregation of Al-Ni-Y for the large negative enthalpy of mixing and decreased the nucleation energy. Besides, deformation induced precipitation of the (Al-Ni-Y)-rich nano phase. At the same time, serious lattice distortion occurred in the nano particle which was shown in the inverse FFT in figures 6(b) and (c). It could be deduced that the second yielding phenomenon was mainly caused by the interaction between nano-precipitates and dislocations. It indicated that the precipitation process of the nano phases could be accelerated by compression and the pinning effect of the precipitates on the dislocations could increase the yielding strength. While, the nano precipitate (marked with gray arrows) were formed around the nano particle in the FCC (L1₂) phase and lattice distortion expanded (marked with black arrows) into the neighboring matrix with the further deformation as shown in figure 6(d).