Dual-phase γ + γ′ microstructure

Ni-based superalloys, as classical materials designed for the high-temperature environment, have been under development for decades since the emergence of rotating engines.^{Reference Sims, Stoloff and Hagel19} There are several reasons for the application of superalloys for industrial applications and especially aircraft applications, including high-temperature strength, creep resistance, and oxidation resistance.^{Reference Sims, Stoloff and Hagel19} The microstructures of some PH-HEAs^{Reference Yang, Zhao, Tong, Jiao, Wei, Cai, Han, Chen, Hu, Kai, Lu, Liu and Liu9,Reference Zhao, Yang, Tong, Wang, Luan, Jiao, Chen, Yang, Hu and Liu21,Reference Gwalani, Choudhuri, Soni, Ren, Styles, Hwang, Nam, Ryu, Hong and Banerjee28–Reference Xu, Liu, Guo, Hirata, Fujita, Nieh, Liu and Chen31} are analogous to those of superalloys, with γ′ precipitates embedded in a γ matrix, (Figure 1a, left and middle), but with high-entropy characteristics both in the matrix and precipitates. Typically, γ′ is a Cu₃Au-type high-entropy intermetallic compound with Ni and Co atoms mainly substituting for the center sites on the cubic faces, and with Al and Ti atoms mainly substituting for the corner sites of the unit cell. Fe and Cr atoms evidently partition into the γ matrix and are depleted in the γ′ precipitation, as shown in Figure 1a, right (black circles represent the precipitates).^{Reference Yang, Zhao, Luan, Han, Wei, Kai and Liu11} Cu has been proven to promote the nucleation of the γ′ phase by providing heterogeneous nucleation sites in the form of Cu clusters,^{Reference Gwalani, Choudhuri, Soni, Ren, Styles, Hwang, Nam, Ryu, Hong and Banerjee28} resulting in a conjoint γ′ phase/Cu-rich cluster within the fcc matrix.

Figure 1.

Nanostructures and microstructures of precipitation-hardened high-entropy alloys (PH-HEAs). (a) (Left) Transmission electron microscope (TEM) image showing the morphology of the L1₂ precipitate. (Inset) Selected-area diffraction patterns taken from zone axis Z = [001]. Red circles show the superlattice diffraction spots of the L1₂-type precipitates, (middle) high-resolution TEM image showing the coherent interface between an L1₂ precipitate and the face-centered-cubic (fcc) matrix, and (right) elemental distribution map obtained by atom probe tomography.^{Reference Yang, Zhao, Luan, Han, Wei, Kai and Liu11} (b) Microstructure of Ni₃₆Co₂₅Fe₁₅Cr₈Al₁₀Ti₆ (at.%) PH-HEAs: (left) high magnification of γ′ precipitation, (middle) dendritic and interdendritic area, and (right) high magnification of needle-like NiAl precipitates (the β phase).^{Reference Daoud, Manzoni, Wanderka and Glatzel32} (c) Detrimental phases: (left) NiAl phase at γ/Laves interfaces during aging in Ni₂₂Co₃₀Fe₁₈Cr₁₈Al₈Nb₄ PH-HEA,^{Reference Antonov, Detrois and Tin34} (middle) needle-like η phase in Ni_1.5Co_1.5FeCrAl_0.1Ti_0.4 PH-HEA, and (right) β phase in Ni_1.5Co_1.5FeCrAl_0.5 PH-HEA.^{Reference Chang and Yeh35} (d) Microstructure of Ni_1.5Co_1.5FeCrAl_0.2Ti_0.3 PH-HEA. (Left) Grain boundary, (middle) magnification of cellular γ′ precipitation along a grain boundary, and (right) well-developed spherical γ′ precipitates within the grain interior.^{Reference Kuo and Tsai37} Note: DP, discontinuous precipitation (cellular γ′ phase); GB, grain boundary.

Detrimental phases

On the path toward optimizing the microstructures of PH-HEAs, the most urgent task is to achieve a pristine γ/γ′ structure without the formation of other undesirable brittle intermetallic phases, especially the TCP phases. Daoud et al.^{Reference Daoud, Manzoni, Wanderka and Glatzel32} and Manzoni et al.^{Reference Manzoni, Singh, Daoud, Popp, Völkl, Glatzel and Wanderka33} observed the formation of a needle-shaped NiAl phase (β phase in Figure 1b, right) and Cr-Fe-rich platelets in Ni₃₆Co₂₅Fe₁₅Cr₈Al₁₀Ti₆ (at.%) and Ni₁₆Co₁₅Fe₁₅Cr₂₃Cu₈Al₂₃ (at.%) PH-HEAs, respectively. Niobium contributed to the formation of the Nb-rich Laves phase during homogenization, while the B2-type NiAl phase was observed at the γ/Laves interfaces during aging in the Ni₂₂Co₃₀Fe₁₈Cr₁₈Al₈Nb₄ (at.%) alloy (Figure 1c, left).^{Reference Antonov, Detrois and Tin34}

Chang et al.^{Reference Chang and Yeh35} found that the needle-like Ni₃Ti phase (the η phase) formed in the Ni_1.5Co_1.5FeCrAl_0.1Ti_0.4 PH-HEA after aging at 1000°C (Figure 1c, middle), while the NiAl phase (the β phase) was observed in the Ni_1.5Co_1.5FeCrAl_0.5 PH-HEA after aging at 900°C (Figure 1c, right). Furthermore, the formation of the η phase was found to be suppressed by replacing Ti with Al addition, but excessive Al will induce the formation of the β phase.^{Reference Chang and Yeh35} The L2₁-type Heusler phase was found in the Ni₃₆Co₂₅Fe₁₅Cr₈Al₁₀Ti₆ (at.%) PH-HEA, which remained even after a long-time annealing at 900°C.^{Reference Haas, Manzoni, Krieg and Glatzel29} The Heusler phase is a Ni₂TiAl-type precipitate with a large unit cell that comprises eight small bcc NiAl-type unit cells. For each small bcc, nickel atoms occupy the corner sites and aluminum or titanium alternatively occupy the body-center site. These undesirable phases are generally hard and brittle, and thus they would act as possible sources for crack initiation and propagation. On the other hand, they consume valuable Al and Ti content, which could be used for elevating the γ′ phase volume fraction and providing alloys with higher strengths at high temperatures.

Cellular precipitation at grain boundaries

Another key issue is the heterogeneous precipitation of cellular γ′ precipitates along grain boundaries,^{Reference Chang and Yeh36,Reference Kuo and Tsai37} which is also observed in traditional alloys, such as Al alloys,^{Reference Williams and Edington38} Cu alloys,^{Reference Zhao and Notis39} and Ni-based superalloys.^{Reference Nystrom, Pollock, Murphy and Garg40} Such discontinuous precipitation behavior^{Reference Kuo and Tsai37} is generally ascribed to the chemical supersaturation in the γ matrix and grain-boundary migration (Figure 1d). It should be noted that tuning the chemical compositions and heat-treatment procedures can avoid such precipitation. Cellular precipitation can possibly be eliminated by changing the quenching method from water quenching to furnace cooling or substituting Co content with Ni. The slower cooling rate is expected to provide more time for γ′ to precipitate at grain boundaries, thus impeding the grain-boundary migration and preventing cellular precipitation. Substituting Co by Ni increases the γ′ solvus temperature, which allows more γ′ to precipitate during the cooling process. The supersaturation degree in the matrix is reduced, resulting in the disappearance of cellular γ′ precipitation. Yeh et al. ^{Reference Chang and Yeh36} and Tsai et al.^{Reference Kuo and Tsai37} proposed that such cellular γ′ precipitates at the grain boundaries have little influence on room-temperature tensile properties, but they cause rapid ductility loss at intermediate temperatures (roughly from 600°C to 800°C), as shown in Figure 2a, which was associated with intergranular fracture.^{Reference Chang and Yeh36,Reference Jensen and Tien41}

Figure 2.

Intermediate temperature embrittlement of precipitation-hardened high-entropy alloys (PH-HEAs) and computational-aided alloy design. (a) (Left) Ductility loss at 800°C in Ni_1.5Co_1.5FeCrAl_0.2Ti_0.3 PH-HEAs.^{Reference Kuo and Tsai37} (Right) Strength and ductility of Ni_27.27Co_27.27Fe_18.18Cr_18.18Al_3.31Ti_5.78 (at.%) PH-HEAs at different temperatures; cellular precipitation caused sharp ductility loss at around 750°C.^{Reference Chang and Yeh36} (b) CALculation of PHase Diagrams (CALPHAD) results of (left) (NiCoCr)_97–xAl_xTi₃ (at.%) PH-HEAs^{Reference Zhao, Yang, Tong, Wang, Luan, Jiao, Chen, Yang, Hu and Liu21} and (right) Ni_1.5Co_1.5FeCrAl_xTi_y (x + y = 0.5) PH-HEAs.^{Reference Chang and Yeh35} (c) Phase diagram predictions of (left) Ni₃₃Co₁₇Fe₁₇Cr₁₇Cu₈Al₈ PH-HEA and (right) Ni_34.8Co₁₇Fe₁₇Cr₁₄Cu₈Al₈Ti₁Mo_0.1W_0.1 (at.%) PH-HEA.^{Reference Manzoni, Daoud, Voelkl, Glatzel and Wanderka50} Arrows in the left panel indicate the expansion of γ′ phase region through small alloy additions of Ti, Mo, and W.

Similarly, the precipitation of brittle grain-boundary oxides and carbides also caused ductility loss with increased temperatures above 600°C in Inconel 718, but ductility at ambient temperatures remained unaffected.^{Reference Rao, Srinivas and Sarma42} Apart from the effects of grain-boundary precipitates,^{Reference Yeh, Chen, Lin, Gan, Chin, Shun, Tsau and Chang1,Reference Cantor, Chang, Knight and Vincent2} there are several other possible interpretations for the embrittlement, such as interstitial impurities segregation,^{Reference Heatherly and George43} environment-induced embrittlement,^{Reference Bricknell and Woodford44} and intrinsic grain-boundary weakness.^{Reference Takasugi, George, Pope and Izumi45} Impurity (such as Bi and S) levels as low as tens of ppm^{Reference Zheng, Chellali, Schlesiger, Baither and Schmitz46,Reference Kraai and Floreen47} are sufficient to induce embrittlement. In one study, pure nickel samples were exposed at 1000°C in air and vacuum before tensile testing in an argon atmosphere. The samples after exposure to air exhibited a higher level of ductility loss than those under vacuum.^{Reference Bricknell and Woodford44} Apart from oxygen, water vapor in air is also thought to dissociate into atomic H on Ni₃Al surfaces, resulting in embrittlement.^{Reference George, Liu and Pope48} Heterogeneous bonding at the grain-boundary area, which causes intrinsic grain-boundary weakness, was considered to be the reason for the fracture surface consisting of the presence of a completely intergranular fracture surface.^{Reference Takasugi, George, Pope and Izumi45}

A fundamental understanding is still lacking at the present time, and more systematic work is required to elucidate the intrinsic mechanisms of the intermediate temperature embrittlement phenomenon.

Computation-aided alloy design

The CALculation of PHase Diagrams (CALPHAD) method, which is based on thermodynamic calculations, offers reliable guidance for alloy design in terms of chemical compositions and heat-treatment conditions. Different from previous trial-and-error routes for alloy design, the CALPHAD method accelerates the search for a target microstructure and composition by computational approaches, which greatly reduces the experiments needed.

It was found that a stable γ + γ′ dual phase can be maintained by adjusting the compositions and heat-treatment temperature to avoid the formation of η and σ phases (Figure 2b, left).^{Reference Zhao, Yang, Tong, Wang, Luan, Jiao, Chen, Yang, Hu and Liu21} He et al. found that the results of the formation of the minor detrimental Heusler phase in (NiCoFeCr)_100–x–yAl_xTi_y PH-HEAs, apart from the γ matrix and γ′ precipitation, were in good agreement with experimental observations.^{Reference He, Wang, Wu, Liu, Mao, Nieh and Lu49} For the Ni_1.5Co_1.5FeCrAl_xTi_y (x + y = 0.5) PH-HEAs, Chang and Yeh^{Reference Chang and Yeh35} successfully predicted the microstructural evolution of (γ + η) in Ti_0.5 to (γ + η + γ′) in Al_0.1Ti_0.4, to (γ + γ′) in Al_0.2Ti_0.3, to (γ + β + γ′) in Al_0.3Ti_0.2 and to (γ + β) in Al_0.5 PH-HEAs (Figure 2b, right). Through a comparison of the phase diagrams of Ni₃₃Co₁₇Fe₁₇Cr₁₇Cu₈Al₈ (at.%) and Ni_34.8Co₁₇Fe₁₇Cr₁₄Cu₈Al₈Ti₁Mo_0.1W_0.1 (at.%) (Figure 2c), it was found that the curve for the γ′ phase was shifted upward, which implied a higher γ′ volume fraction in the Ni_34.8Co₁₇Fe₁₇Cr₁₄Cu₈Al₈Ti₁Mo_0.1W_0.1 (at.%) PH-HEA than the Ni₃₃Co₁₇Fe₁₇Cr₁₇Cu₈Al₈ (at.%) PH-HEA.^{Reference Manzoni, Daoud, Voelkl, Glatzel and Wanderka50} It is of interest to point out that small additions of refractory elements Mo, W, and Ti can raise the volume fraction of the γ′ phase considerably, resulting in the suppression of the NiAl.^{Reference Manzoni, Daoud, Voelkl, Glatzel and Wanderka50}

However, some discrepancies still exist between the experimental results and the CALPHAD predictions^{Reference Liu, Lu, He, Luan, Wang, Liu, Liu, Chen and Liu13,Reference Tsao, Yeh, Kuo and Murakami51} in terms of detrimental phases, volume fractions of γ′, and γ′ solvus temperatures. It appears that more thermodynamic data for CALPHAD calculations are needed for PH-HEAs design.

Summary

The microstructural evolution of PH-HEAs has been reviewed in this article. A dual high-entropy nanostructure consisting of high-entropy γ′ precipitates and a high-entropy γ matrix is analogous to that of conventional superalloys. Due to severe lattice distortion, sluggish diffusion, and the high-entropy effect, PH-HEAs have the potential to replace conventional Ni-based superalloys and to be used in high-temperature environments. Some critical issues remain, suggesting future directions for research:

1. 1. Phase control should be given higher priority. Hard and brittle deleterious phases need to be eliminated by carefully tuning chemical compositions and heat treatments. The CALPHAD method makes it easier to obtain a target microstructure without time-consuming experiments. Retaining a stable γ′/γ dual-phase structure at higher temperatures is a presumption for the design of high-temperature structural materials. Phase stability evaluation is an important consideration in selecting structural materials under elevated-temperature environments. Besides, a high-volume fraction of the γ′ strengthener is preferred in order to develop high-performance structural materials. The preferential size, morphological shape, and composition of the precipitates, as well as lattice misfit between the matrix and the precipitates in PH-HEAs for high-temperature applications, remain open questions for researchers.

2. 2. The high-temperature strength of PH-HEAs decreased rapidly when the temperature approached 1000°C because of the rapid dissolution of hardening precipitates. The yield strength of the Ni₃₆Co₂₅Fe₁₅Cr₈Al₁₀Ti₆ PH-HEA remained at 672 MPa at 800°C, but dropped to just 148 MPa at 1000°C.^{Reference Daoud, Manzoni, Wanderka and Glatzel32} Similarly, the hardnesses of Ni_40.7Co_20.6Fe_11.5Cr_12.2Al_7.8Ti_7.2 (at.%) and Ni_48.6Co_17.0Fe_9.0Cr_7.5Al_10.3Ti_5.8Ta_0.6Mo_0.8W_0.4 (at.%) PH-HEAs were superior to that of Ni-based superalloy CM247LC (Ni-10Co-8.4Cr-5.5Al-1.0Ti-3.0Ta-0.7Mo-10.0W-0.15C-0.015B-0.05Zr-1.5Hf wt%) at room temperature, but softened quickly when the temperature increased above 600°C.^{Reference Tsao, Yeh, Kuo and Murakami24} The stabilization of the L1₂ precipitate at high temperatures is a key issue for alloy design.

3. 3. There is limited data on creep resistance^{Reference Tsao, Yeh, Kuo, Kakehi, Murakami, Yeh and Jian30} and thermal fatigue behavior of PH-HEAs at the present time. Also, the mechanical performance of directionally solidified (DS) and single-crystal (SC) PH-HEAs have been scarcely reported.^{Reference Tsao, Yeh, Kuo, Kakehi, Murakami, Yeh and Jian30,Reference Manzoni, Haas, Daoud, Glatzel, Förster and Wanderka52} There are three main reasons for the enhancement of elevated-temperature properties in DS- and SC PH-HEAs. The elimination of grain boundaries normal to the stress axis improves creep resistance and high-temperature ductility because of the lack of failure initiation sites. The second is that a preferred low-modulus [001] orientation in the solidification direction improves thermal fatigue resistance. Third, due to the removal of grain boundaries along the stress axis in the DS PH-HEAs and the complete removal of grain boundaries in the SC PH-HEAs, incipient melting temperatures for such PH-HEAs are higher than their polycrystal counterparts, implying a better high-temperature performance. More importantly, we need to compare the high-temperature strength and creep resistance of DS- and SC PH-HEAs with that of Ni-based superalloys.

4. 4. High-temperature oxidation resistance is another key issue to be characterized. Hot corrosion, as a form of degradation, should also be considered. PH-HEAs mainly count on Al and Cr content to form alumina and chromia, respectively, for surface protection and oxidation resistance. As a result, the alloying additions of Al and Cr should be quantitively studied in order to improve the PH-HEA oxidation and corrosion resistance.