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Chen-Yuan Wang, [Sae Matsunaga](https://orcid.org/0000-0001-7285-5372), [Yoshiaki Toda](https://orcid.org/0000-0002-8343-2890), [Hideyuki Murakami](https://orcid.org/0000-0001-8220-5816), [An-Chou Yeh](https://orcid.org/0000-0002-9460-8345), [Yoko Yamabe-Mitarai](https://orcid.org/0000-0002-9674-5503)

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[Effect of Alloying Elements on the High-Temperature Yielding Behavior of Multicomponent γ′-L12 Alloys](https://mdr.nims.go.jp/datasets/ceeaceca-8c43-44d7-aa61-a828c28dbbb1)

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Effect of Alloying Elements on the High-Temperature Yielding Behavior of Multicomponent '-L12 AlloysCitation: Wang, C.-Y.; Matsunaga, S.;Toda, Y.; Murakami, H.; Yeh, A.-C.;Yamabe-Mitarai, Y. Effect of AlloyingElements on the High-TemperatureYielding Behavior of Multicomponentγ′-L12 Alloys. Materials 2024, 17, 2280.https://doi.org/10.3390/ma17102280Academic Editor: Joan-Josep SuñolReceived: 19 April 2024Revised: 3 May 2024Accepted: 9 May 2024Published: 11 May 2024Copyright: © 2024 by the authors.Licensee MDPI, Basel, Switzerland.This article is an open access articledistributed under the terms andconditions of the Creative CommonsAttribution (CC BY) license (https://creativecommons.org/licenses/by/4.0/).materialsArticleEffect of Alloying Elements on the High-Temperature YieldingBehavior of Multicomponent γ′-L12 AlloysChen-Yuan Wang 1, Sae Matsunaga 1,* , Yoshiaki Toda 2 , Hideyuki Murakami 2,3 , An-Chou Yeh 4and Yoko Yamabe-Mitarai 1,*1 Department of Advanced Materials Science, The University of Tokyo, 5-1-5 Kashiwanoha,Kashiwa-shi 277-8561, Chiba, Japan; 1775090211@edu.k.u-tokyo.ac.jp2 National Institute for Materials Science, 1-2-1 Sengen, Tsukuba 305-0047, Ibaraki, Japan;toda.yoshiaki@nims.go.jp (Y.T.); murakami.hideyuki@nims.go.jp (H.M.)3 Department of Nanoengineering and Nanoscience, Waseda University, Shinjuku 169-8555, Tokyo, Japan4 High Entropy Materials Center, National Tsing Hua University, Hsinchu 30013, Taiwan;yehac@mx.nthu.edu.tw* Correspondence: smatsun@edu.k.u-tokyo.ac.jp (S.M.); mitarai.yoko@edu.k.u-tokyo.ac.jp (Y.Y.-M.)Abstract: The exceptional mechanical properties of Ni-based high entropy alloys are due to thepresence of ordered L12 (γ′) precipitates embedded within a disordered matrix phase. While thestrengthening contribution of the γ′ phase is generally accepted, there is no consensus on the precisecontribution of the individual strengthening mechanisms to the overall strength. In addition, changesin alloy composition influence several different mechanisms, making the assessment of alloyingconditions complex. Multicomponent L12-ordered single-phase alloys were systematically developedwith the aid of CALPHAD thermodynamic calculations. The alloying elements Co, Cr, Ti, and Nbwere chosen to complexify the Ni3Al structure. The existence of the γ′ single phase was validatedby microstructure characterization and phase identification. A high-temperature compression testfrom 500 ◦C to 1000 ◦C revealed a positive temperature dependence of strength before reaching thepeak strength in the studied alloys NiCoCrAl, NiCoCrAlTi, and NiCoCrAlNb. Ti and Nb alloyingaddition significantly enhanced the high-temperature yield strengths before the peak temperature.The yield strength was modeled by summing the individual effects of solid solution strengthening,grain boundary strengthening, order strengthening, and cross-slip-induced strengthening. Cross-slip-induced strengthening was shown to be the key contributor to the high-temperature strengthenhancement.Keywords: multicomponent alloy; high-temperature deformation; intermetallic compounds; CAL-PHAD1. IntroductionNi-based superalloys have been extensively employed as critical high-temperaturestructural materials used in aerospace engineering and the power generation industrydue to their superior high-temperature mechanical properties, which are derived fromthe combination of 60–70% volume fraction of ordered γ′ precipitates (L12) coherentlyembedded in the solid-solution γ matrix (FCC-A1) [1]. The alloy design of Ni-based su-peralloys has approached its limit for performance enhancements due to the constraintsof Ni-based systems. A novel concept in alloy design, high entropy alloys (HEAs), hasmade a breakthrough in conventional alloys by revisiting the vast composition space [2].Precipitation-strengthened HEAs, as newly emerging structural materials, are highly attrac-tive due to their combination of strength and ductility [3–9]. Compared to grain boundarystrengthening and solid solution strengthening, precipitation strengthening of multicom-ponent L12 nanoprecipitates can significantly improve the yield strength of HEAs whilemaintaining good ductility. Yeh et al. [9] first introduced the L12 nanoprecipitates of (Ni,Materials 2024, 17, 2280. https://doi.org/10.3390/ma17102280 https://www.mdpi.com/journal/materialshttps://doi.org/10.3390/ma17102280https://doi.org/10.3390/ma17102280https://creativecommons.org/https://creativecommons.org/licenses/by/4.0/https://creativecommons.org/licenses/by/4.0/https://www.mdpi.com/journal/materialshttps://www.mdpi.comhttps://orcid.org/0000-0001-7285-5372https://orcid.org/0000-0002-8343-2890https://orcid.org/0000-0001-8220-5816https://orcid.org/0000-0002-9460-8345https://orcid.org/0000-0002-9674-5503https://doi.org/10.3390/ma17102280https://www.mdpi.com/journal/materialshttps://www.mdpi.com/article/10.3390/ma17102280?type=check_update&version=2Materials 2024, 17, 2280 2 of 12Co)3Ti into the FeCoNiCr system by adding Ti to form a precipitation-strengthened HEA.Later, Yang et al. [8] designed and developed a complex alloy (FeNiCo)86Al7Ti7 with multi-component L12 intermetallic particles (Ni, Co, Fe)3(Ti, Al, Fe) and reported a combinationof high strength (~1 GPa) and excellent ductility (~50%) at room temperature. They at-tributed the pronounced increase in yield strength to the high ordering strengthening bythe high density of L12 nanoprecipitates and the higher anti-phase boundary energy (APBE)provided by Ti addition.Recently, Yang et al. [10] developed a multicomponent L12-ordered alloy, NiCoFeAlTiB,and observed that the materials did not experience obvious softening behavior below800 ◦C under the hardness test. Long et al. [11] investigated a multicomponent Co-basedL12 ordered single phase intermetallic alloy and found that the alloy exhibited anomalousyield strength increase from 250 ◦C to 800 ◦C. The stress anomaly was stronger thanternary L12 Co3(Al, W) and B-doped Ni3Al. These results revealed that a high-temperaturecapability could be achieved by the complex L12 intermetallic. It is well known that the L12phase exhibits anomalous yielding behavior, represented by a substantial yield strengthincrease at intermediate temperatures [1]. This positive temperature dependence of stressof L12 precipitates confers excellent high-temperature mechanical properties. It has beengenerally accepted that the anomalous yielding behavior was due to the Kear–Wilsdorfmechanism, forming anti-phase boundary (APB)-coupled dislocation pairs by thermallyactivated cross-slip of a screw segment of the superpartial dislocation from octahedral {111}planes to cube {010} planes. While the strength in Ni-based superalloys is proportional tothe APBE, the anisotropy in APBE between octahedral {111} planes and cube {010} planesdetermines the yield strength anomaly. Since a perfect dislocation slipping on the plane inNi3Al consists of four Shockley partial dislocations and bounding one APB fault and twocomplex stacking faults (CSF), the occurrence of thermally activated cross-slip process isdependent on APB and CSF energies [12,13].Considerable efforts have been devoted to researching the composition dependenceof the high-temperature strength and associated planar fault energies [14–18]. Mishimaet al. [14,19] systematically investigated the effect of ternary addition on the temperaturedependence of strength in ternary Ni3Al compounds. Among the investigated alloyingelements, Ti, Zr, Hf, V, Nb, Ta, Mo, W, Ti, Nb, and Ta produced the most drastic strengthincrement. Moreover, advances in computational models and modern computers allow forthe APBE of the ternary Ni3Al intermetallic compounds to be predicted by density functiontheory (DFT) calculations [13,20–24]. Gorbatov et al. [23] studied the effect of compositionon the APBE on both octahedral planes and cube planes of Ni-based L12-ordered alloys byab initio calculations and reported ternary alloying additions. Ti, V, and Cr occupied the Alsublattice site and increased the APBE for both {111} and {010} planes, while Co, Cu, andFe occupied the Ni sublattice site and sightly affected both the {111} and {010} APBE.Recent studies on the deformation mechanisms of L12-strengthened alloys revealedthat γ′-strengtheners, such as Ti, Ta, and Nb, significantly improved the creep performanceby forming the other phases at planar defects, which are considered to have slower de-formation kinetics [25]. In addition, the experimental and computational works showedCo and Cr and γ′ strengtheners significantly affected the phase transformation, whichpreferentially occurred at planar defects. While forming an ordered phase with slowerformation kinetics can enhance the high-temperature creep performance, the formationof disordered phases will degrade it. Determining the resulting phase hinges on subtlediscrepancies in nominal alloy composition [25–28]. Although the deformation mechanismsof γ′-strengthened alloys have been investigated for decades, there are still some uncoveredquestions about the γ′ phase. In γ + γ′ structures, deformation mechanisms are usuallysignificantly affected by the interactions at γ and γ′ interface [29]; thus, elucidating thestrengthening effect solely attributable to the γ′ phase has proven to be a formidable task.Furthermore, while numerous recent studies have been dedicated to the effect of alloyingelements on the γ′-shearing mechanisms, their influence on phenomena unique to the L12Materials 2024, 17, 2280 3 of 12structure, such as the elevation of yielding stress at high temperatures, remains significantlyunderexplored in a systematic manner.Fundamental and systematic studies on multicomponent L12 phases are urgentlyrequired to accelerate the advancement of L12-strengthened alloys. The precipitationstrengthening by L12-type Ni3Al nanoparticles has been extensively investigated in Ni-based superalloys and HEAs. However, most previously investigated Ni3Al phases arecompositionally simple and contain a lower level of ternary elements. Because of theextensive alloying addition of Cr and Co in Ni-based superalloys and L12 precipitationstrengthened HEAs and their effect on local phase transformation, these two elementswere chosen as a base with Ni3Al to design multicomponent L12 alloys. This work sys-tematically varied compositions by substituting Ni and Al with Ti or Nb to study thecompositional effects on the yielding behaviors of the L12 phase. The investigated multi-component L12 alloys were systematically designed using the thermodynamics softwareCALPHAD Thermo-Calc TCNI8 Ni-based Superalloys Database [30]. The yielding behav-iors of the multicomponent L12 alloys were evaluated by high-temperature compressiontests. The analysis and discussion will delve into the contribution of different strengtheningmechanisms under various alloying conditions.2. Materials and MethodsThe multicomponent γ′ (L12-ordered) single phase was designed by utilizing theThermo-Calc© TCNI8 database [30], aiming to obtain L12-ordered single phase within thetemperature range between 800 ◦C and 950 ◦C. We developed three systems, as follows:• NiCoCrAl,• NiCoCrAlTi,• NiCoCrAlNb.The pseudo-binary isopleths and volume fraction diagrams of three multicomponentsystems were calculated to confirm the solubility limit of Ti or Nb alloying addition inNiCoCrAl alloy without forming any secondary phases. Ti and Nb were carefully addedto avoid exceeding their solubility limits, approximately 2 at%, in the NiCoCrAl alloy.Selected compositions for the studied alloys are listed in Table 1 and marked using thedash lines in Figure 1.Table 1. Nominal compositions of the studied alloys (at%).Al Co Cr Ni Ti NbCY1 21 10 5 64 - -CY2 20 10 5 63 2 -CY3 20 10 5 63 - 2Materials 2024, 17, x FOR PEER REVIEW 4 of 12    Figure 1. The pseudo-binary isopleth of three different multicomponent systems: (a) NiCo10Cr5Al, (b) NiCo10Cr5Al20Ti, and (c) NiCo10Cr5Al20Nb.  Table 1. Nominal compositions of the studied alloys (at%).  Al Co Cr Ni Ti Nb CY1 21 10 5 64 - - CY2 20 10 5 63 2 - CY3 20 10 5 63 - 2 3. Results and Discussion 3.1. Microstructures The volume fraction diagrams of the designed alloys are shown in Figure 2. The γ’ (L12-ordered) single phase exists in the range of 659–1104 °C in CY1, 870–1060 °C in CY2, and 858–991 °C in CY3. The temperature window with only the L12-ordered single phase narrowed by adding the fifth element, while the γ + γ’ two-phase window broadened. The γ + γ′ window is 1104–1163 °C in CY1, 1060–1171 °C in CY2, and 991–1133 °C in CY3. Both Ti and Nb were classified as γ′-stabilizers in terms of the alloy design strategy for γ + γ′ two-phase Ni-based superalloys [1]. Our calculation results indicated that these elements stabilized the L12-ordered phase in a wide range of temperatures with the existence of the γ matrix while they decreased the L12-ordered single-phase window. We also investigated the site occupancy in each alloy using thermodynamic calculations. It was simulated that Co tended to occupy the Ni sublattice sites, while Cr, Ti, and Nb atoms substituted the Al site atoms to create (Ni, Co)3(Al, Cr) in CY1, (Ni, Co)3(Al, Cr, Ti) in CY2, and (Ni, Co)3(Al, Cr, Nb) in CY3.  Figure 2. Volume fraction diagrams of the following multicomponent L12-ordered intermetallic compounds: (a) CY1 (NiCoCrAl), (b) CY2 (NiCoCrAlTi), and (c) CY3 (NiCoCrAlNb).  Figure 1. The pseudo-binary isopleth of three different multicomponent systems: (a) NiCo10Cr5Al,(b) NiCo10Cr5Al20Ti, and (c) NiCo10Cr5Al20Nb.Materials 2024, 17, 2280 4 of 12In total, 300 g of the studied alloy ingots were prepared by the vacuum arc melt-ing (VAM) process under an argon atmosphere with a Ti getter. Repeated melting wascarried out at least six times to ensure the chemical homogeneity. As-cast samples werehomogenized at 980 ◦C for 168 h, followed by air cooling to room temperature. For CY3alloy with Nb alloying addition, a prior solution heat treatment (SHT) at 1280 ◦C for 168 hwas necessary to eliminate the dendritic structure caused by the severe Nb segregation.The constituent phases were identified using an X-ray diffractometer (XRD, SmartLab,Rigaku, Tokyo, Japan). The microstructures of each alloy were characterized using scan-ning electron microscopy (SEM, JSM-7200 F, JEOL, Tokyo, Japan). The phase transformationtemperatures of the studied alloys were determined by differential thermal analysis (DTA,Labsys DSC/DTA, Setaram, Caluire-et-Cuire, France) experiments between 25 ◦C and1500 ◦C, with a scanning rate of 10 ◦C/min. Cuboid compression specimens with 3 mm and6 mm width and height length were sectioned by a high-speed precision cutting machine.High-temperature compression tests were performed to understand the strength evolutionat 500, 600, 700, 800, 900, and 1000 ◦C with a strain rate of 10−4 s−1. All the tests wereinterrupted after reaching approximately 10% plastic strain.3. Results and Discussion3.1. MicrostructuresThe volume fraction diagrams of the designed alloys are shown in Figure 2. The γ’(L12-ordered) single phase exists in the range of 659–1104 ◦C in CY1, 870–1060 ◦C in CY2,and 858–991 ◦C in CY3. The temperature window with only the L12-ordered single phasenarrowed by adding the fifth element, while the γ + γ’ two-phase window broadened. Theγ + γ′ window is 1104–1163 ◦C in CY1, 1060–1171 ◦C in CY2, and 991–1133 ◦C in CY3. BothTi and Nb were classified as γ′-stabilizers in terms of the alloy design strategy for γ + γ′two-phase Ni-based superalloys [1]. Our calculation results indicated that these elementsstabilized the L12-ordered phase in a wide range of temperatures with the existence of theγ matrix while they decreased the L12-ordered single-phase window. We also investigatedthe site occupancy in each alloy using thermodynamic calculations. It was simulated thatCo tended to occupy the Ni sublattice sites, while Cr, Ti, and Nb atoms substituted the Alsite atoms to create (Ni, Co)3(Al, Cr) in CY1, (Ni, Co)3(Al, Cr, Ti) in CY2, and (Ni, Co)3(Al,Cr, Nb) in CY3.Materials 2024, 17, x FOR PEER REVIEW 4 of 12    Figure 1. The pseudo-binary isopleth of three different multicomponent systems: (a) NiCo10Cr5Al, (b) NiCo10Cr5Al20Ti, and (c) NiCo10Cr5Al20Nb.  Table 1. Nominal compositions of the studied alloys (at%).  Al Co Cr Ni Ti Nb CY1 21 10 5 64 - - CY2 20 10 5 63 2 - CY3 20 10 5 63 - 2 3. Results and Discussion 3.1. Microstructures The volume fraction diagrams of the designed alloys are shown in Figure 2. The γ’ (L12-ordered) single phase exists in the range of 659–1104 °C in CY1, 870–1060 °C in CY2, and 858–991 °C in CY3. The temperature window with only the L12-ordered single phase narrowed by adding the fifth element, while the γ + γ’ two-phase window broadened. The γ + γ′ window is 1104–1163 °C in CY1, 1060–1171 °C in CY2, and 991–1133 °C in CY3. Both Ti and Nb were classified as γ′-stabilizers in terms of the alloy design strategy for γ + γ′ two-phase Ni-based superalloys [1]. Our calculation results indicated that these elements stabilized the L12-ordered phase in a wide range of temperatures with the existence of the γ matrix while they decreased the L12-ordered single-phase window. We also investigated the site occupancy in each alloy using thermodynamic calculations. It was simulated that Co tended to occupy the Ni sublattice sites, while Cr, Ti, and Nb atoms substituted the Al site atoms to create (Ni, Co)3(Al, Cr) in CY1, (Ni, Co)3(Al, Cr, Ti) in CY2, and (Ni, Co)3(Al, Cr, Nb) in CY3.  Figure 2. Volume fraction diagrams of the following multicomponent L12-ordered intermetallic compounds: (a) CY1 (NiCoCrAl), (b) CY2 (NiCoCrAlTi), and (c) CY3 (NiCoCrAlNb).  Figure 2. Volume fraction diagrams of the following multicomponent L12-ordered intermetalliccompounds: (a) CY1 (NiCoCrAl), (b) CY2 (NiCoCrAlTi), and (c) CY3 (NiCoCrAlNb).Figure 3 shows the microstructure of the CY1, CY2, and CY3 alloys after homogenizingheat treatment. All alloys exhibited single-phase microstructures after the heat treatments.The average grain size of each alloy is around 503 µm in CY1, 849 µm in CY2, and 323 µmin CY3. Any secondary phases were not detected either within the grains or at the grainboundaries, as shown in Figure 3d–f, which confirms that we successfully designed andmanufactured the multicomponent L12-ordered single-phase alloys.Materials 2024, 17, 2280 5 of 12Materials 2024, 17, x FOR PEER REVIEW 5 of 12   Figure 3 shows the microstructure of the CY1, CY2, and CY3 alloys after homogeniz-ing heat treatment. All alloys exhibited single-phase microstructures after the heat treat-ments. The average grain size of each alloy is around 503 µm in CY1, 849 µm in CY2, and 323 µm in CY3. Any secondary phases were not detected either within the grains or at the grain boundaries, as shown in Figure 3d–f, which confirms that we successfully designed and manufactured the multicomponent L12-ordered single-phase alloys.  Figure 3. Backscatter images of microstructures after the heat treatments under different magnifica-tions: (a) CY1 (250×), (b) CY2 (250×), (c) CY3 (250×), (d) CY1 (10,000×), (e) CY2 (10,000×), and (f) CY3 (10,000×). Figure 4 shows the indexed XRD spectra peaks. The 2θ values, corresponding to the indexed diffraction peaks of the studied alloys, are listed with the ones of Ni3Al in JCPDS files in Table 2. The 2θ values of the studied alloys CY1, CY2, and CY3 agreed well with the values in JCPDS files, which confirmed the existence of an L12-ordered single phase in the studied alloys. Minor differences from the JCPDS files could arise from the peak broadening due to differences in grain size, lattice strain, and changes in lattice parameters [31]. The crystal structure and microstructure of the alloys after the heat treatments were in good agreement with the prediction from the calculated equilibrium phase diagrams.  Figure 4. XRD analysis of the multicomponent L12-ordered intermetallic alloys. Figure 3. Backscatter images of microstructures after the heat treatments under different magnifi-cations: (a) CY1 (250×), (b) CY2 (250×), (c) CY3 (250×), (d) CY1 (10,000×), (e) CY2 (10,000×), and(f) CY3 (10,000×).Figure 4 shows the indexed XRD spectra peaks. The 2θ values, corresponding to theindexed diffraction peaks of the studied alloys, are listed with the ones of Ni3Al in JCPDSfiles in Table 2. The 2θ values of the studied alloys CY1, CY2, and CY3 agreed well with thevalues in JCPDS files, which confirmed the existence of an L12-ordered single phase in thestudied alloys. Minor differences from the JCPDS files could arise from the peak broadeningdue to differences in grain size, lattice strain, and changes in lattice parameters [31]. Thecrystal structure and microstructure of the alloys after the heat treatments were in goodagreement with the prediction from the calculated equilibrium phase diagrams.Materials 2024, 17, x FOR PEER REVIEW 5 of 12   Figure 3 shows the microstructure of the CY1, CY2, and CY3 alloys after homogeniz-ing heat treatment. All alloys exhibited single-phase microstructures after the heat treat-ments. The average grain size of each alloy is around 503 µm in CY1, 849 µm in CY2, and 323 µm in CY3. Any secondary phases were not detected either within the grains or at the grain boundaries, as shown in Figure 3d–f, which confirms that we successfully designed and manufactured the multicomponent L12-ordered single-phase alloys.  Figure 3. Backscatter images of microstructures after the heat treatments under different magnifica-tions: (a) CY1 (250×), (b) CY2 (250×), (c) CY3 (250×), (d) CY1 (10,000×), (e) CY2 (10,000×), and (f) CY3 (10,000×). Figure 4 shows the indexed XRD spectra peaks. The 2θ values, corresponding to the indexed diffraction peaks of the studied alloys, are listed with the ones of Ni3Al in JCPDS files in Table 2. The 2θ values of the studied alloys CY1, CY2, and CY3 agreed well with the values in JCPDS files, which confirmed the existence of an L12-ordered single phase in the studied alloys. Minor differences from the JCPDS files could arise from the peak broadening due to differences in grain size, lattice strain, and changes in lattice parameters [31]. The crystal structure and microstructure of the alloys after the heat treatments were in good agreement with the prediction from the calculated equilibrium phase diagrams.  Figure 4. XRD analysis of the multicomponent L12-ordered intermetallic alloys. Figure 4. XRD analysis of the multicomponent L12-ordered intermetallic alloys.Table 2. The 2θ (◦) values with related diffraction plane of L12-ordered phase and the correspondingvalues in JCPDS files and the lattice constants of the studied alloys.(100) (111) (200) (220) (311) Lattice Constant(Å)JCPDS 24.90 43.60 50.70 75.02 91.21 -CY1 24.98 43.87 51.51 75.02 91.50 3.48CY2 24.78 43.98 50.59 74.81 91.14 3.58CY3 24.77 44.03 51.00 75.24 91.18 3.59Materials 2024, 17, 2280 6 of 12Figure 5 shows the DTA heating curves of the heat-treated CY1, CY2, and CY3 alloys.The onset temperature of endothermic peaks was determined as the phase transformationtemperatures and solvus temperatures. The phase transformation temperatures (PTTs)were defined as the temperatures at which the alloys could retain the L12-ordered singlephase. Above these temperatures, the secondary and third phases started to form. ThePTTs and solvus temperatures are summarized in Table 3. The PTTs of the CY1, CY2 andCY3 alloys were 1217 ◦C, 1244 ◦C, and 1298 ◦C, respectively. The results showed thatthe L12-ordered single phase persisted at higher temperatures in the quinary alloys withTi or Nb alloying additions. In the Ni-based superalloys with γ + γ′ structure, addingTi or Nb increased γ′-solvus temperatures, indicating these elements could stabilize theL12-ordered phase at elevated temperatures. Since our alloys were designed only to havean L12-ordered single phase, the aforementioned effect of γ′-formers was observed as theexpansion of an L12-ordered single-phase regime. The solidus temperatures for the CY1,CY2, and CY3 alloys were 1353 ◦C, 1324 ◦C, and 1305 ◦C, respectively. The decrease insolidus caused by Ti alloying addition was also observed in Ni-based superalloys due tothe strong partition tendency of Ti to liquid phase, forming segregation phases with lowermelting temperatures and reducing the solidus temperature of the alloy [32,33].Materials 2024, 17, x FOR PEER REVIEW 6 of 12   Table 2. The 2θ (°) values with related diffraction plane of L12-ordered phase and the corresponding values in JCPDS files and the lattice constants of the studied alloys.  (100) (111) (200) (220) (311) Lattice Constant (Å) JCPDS 24.90 43.60 50.70 75.02 91.21 - CY1 24.98 43.87 51.51 75.02 91.50 3.48 CY2 24.78 43.98 50.59 74.81 91.14 3.58 CY3 24.77 44.03 51.00 75.24 91.18 3.59 Figure 5 shows the DTA heating curves of the heat-treated CY1, CY2, and CY3 alloys. The onset temperature of endothermic peaks was determined as the phase transformation temperatures and solvus temperatures. The phase transformation temperatures (PTTs) were defined as the temperatures at which the alloys could retain the L12-ordered single phase. Above these temperatures, the secondary and third phases started to form. The PTTs and solvus temperatures are summarized in Table 3. The PTTs of the CY1, CY2 and CY3 alloys were 1217 °C, 1244 °C, and 1298 °C, respectively. The results showed that the L12-ordered single phase persisted at higher temperatures in the quinary alloys with Ti or Nb alloying additions. In the Ni-based superalloys with γ + γ′ structure, adding Ti or Nb increased γ′-solvus temperatures, indicating these elements could stabilize the L12-or-dered phase at elevated temperatures. Since our alloys were designed only to have an L12-ordered single phase, the aforementioned effect of γ′-formers was observed as the expan-sion of an L12-ordered single-phase regime. The solidus temperatures for the CY1, CY2, and CY3 alloys were 1353 °C, 1324 °C, and 1305 °C, respectively. The decrease in solidus caused by Ti alloying addition was also observed in Ni-based superalloys due to the strong partition tendency of Ti to liquid phase, forming segregation phases with lower melting temperatures and reducing the solidus temperature of the alloy [32,33].  Figure 5. DTA heating curves of the studied alloys: (a) CY1, (b) CY2, and (c) CY3. Table 3. Phase transformation temperatures and solidus temperatures for the studied alloys.  CY1 (°C) CY2 (°C) CY3 (°C) Phase transformation temperature 1217 1244 1298 Solidus temperature 1353 1324 1305 3.2. Yielding Behavior through Temperatures Figure 6 depicts the 0.2% yield stress measured from the high-temperature compres-sion tests for CY1, CY2, and CY3 along the deformation temperatures. The yield stress measured at each temperature is summarized in Table 4. Regardless of the testing temper-ature, CY3 exhibited the highest yield stress among the three alloys before the peak Figure 5. DTA heating curves of the studied alloys: (a) CY1, (b) CY2, and (c) CY3.Table 3. Phase transformation temperatures and solidus temperatures for the studied alloys.CY1 (◦C) CY2 (◦C) CY3 (◦C)Phase transformation temperature 1217 1244 1298Solidus temperature 1353 1324 13053.2. Yielding Behavior through TemperaturesFigure 6 depicts the 0.2% yield stress measured from the high-temperature compres-sion tests for CY1, CY2, and CY3 along the deformation temperatures. The yield stressmeasured at each temperature is summarized in Table 4. Regardless of the testing tem-perature, CY3 exhibited the highest yield stress among the three alloys before the peaktemperature. All alloys exhibited yield stress increasing monotonically before reaching amaximum and then significantly decreasing with increasing the testing temperature afterreaching the peak temperature. The difference in the yield stresses significantly decreasedafter the peak temperatures.Table 4. Yield strengths of the studied alloys at different temperatures.(◦C) CY1 (MPa) CY2 (MPa) CY3 (MPa)500 324.7 420.5 587.7600 391.2 510.6 655.9700 426.5 677.7 690.4800 511.5 617.6 685.2900 417.1 407.0 433.71000 259.5 300.9 288.7Materials 2024, 17, 2280 7 of 12Materials 2024, 17, x FOR PEER REVIEW 7 of 12   temperature. All alloys exhibited yield stress increasing monotonically before reaching a maximum and then significantly decreasing with increasing the testing temperature after reaching the peak temperature. The difference in the yield stresses significantly decreased after the peak temperatures. Comparing the yield stress of CY1, the ones of CY2 are 30% higher at 500 °C, 31% higher at 600 °C, and 59% higher at 700 °C. In addition, the yield stresses of CY3 are 81% higher at 500 °C, 68% higher at 600 °C, and 62% higher at 700 °C. The peak strengths for CY1, CY2, and CY3 alloys were 426 MPa, 678 MPa, and 690 MPa, respectively. Our results revealed that the fifth alloying addition produced a significant enhancement in the yield stress while exhibiting anomalous yielding behavior, which was reported in numerous studies about the high-temperature mechanical behavior of L12-Ni3Al. Mishima et al. [14] investigated the effect of ternary addition of several fourth, fifth, and sixth group elements on the high-temperature mechanical response of polycrystalline Ni3Al and reported that 4 at% Ti addition, which substituted at Al sites, reached the peak strength of about 560 MPa and 2 at% Nb substitution obtained the peak strength of about 620 MPa at 600 °C. It has been generally accepted that Co and Cr addition in Ni-based superalloys strengthen the γ matrix [1]. However, in this research, comparing the compression results of poly-crystalline Ni3Al from Lopez and Hancock [17] with CY2 and CY3, a noticeable strength enhancement was obtained by adding Co and Cr in Ni3Al. Furthermore, comparing the three alloys studied in the research, the Ti or Nb alloying addition showed a significant increase in peak strength.  Figure 6. Temperature dependence of strengths. Table 4. Yield strengths of the studied alloys at different temperatures. (°C) CY1 (MPa) CY2 (MPa) CY3 (MPa) 500  324.7 420.5 587.7 600  391.2 510.6 655.9 700  426.5 677.7 690.4 800  511.5 617.6 685.2 900  417.1 407.0 433.7 1000  259.5 300.9 288.7 Figure 6. Temperature dependence of strengths.Comparing the yield stress of CY1, the ones of CY2 are 30% higher at 500 ◦C, 31%higher at 600 ◦C, and 59% higher at 700 ◦C. In addition, the yield stresses of CY3 are 81%higher at 500 ◦C, 68% higher at 600 ◦C, and 62% higher at 700 ◦C. The peak strengths forCY1, CY2, and CY3 alloys were 426 MPa, 678 MPa, and 690 MPa, respectively. Our resultsrevealed that the fifth alloying addition produced a significant enhancement in the yieldstress while exhibiting anomalous yielding behavior, which was reported in numerousstudies about the high-temperature mechanical behavior of L12-Ni3Al. Mishima et al. [14]investigated the effect of ternary addition of several fourth, fifth, and sixth group elementson the high-temperature mechanical response of polycrystalline Ni3Al and reported that 4at% Ti addition, which substituted at Al sites, reached the peak strength of about 560 MPaand 2 at% Nb substitution obtained the peak strength of about 620 MPa at 600 ◦C. It hasbeen generally accepted that Co and Cr addition in Ni-based superalloys strengthen the γmatrix [1]. However, in this research, comparing the compression results of polycrystallineNi3Al from Lopez and Hancock [17] with CY2 and CY3, a noticeable strength enhancementwas obtained by adding Co and Cr in Ni3Al. Furthermore, comparing the three alloysstudied in the research, the Ti or Nb alloying addition showed a significant increase inpeak strength.During the high-temperature compression deformation up to the peak temperature,the yield strength of the alloys depends on the solid solution strengthening (∆σSS), grainboundary strengthening by the Hall–Petch effect (∆σgb), ordering strengthening (∆σOS),and cross-slip-induced strengthening (∆σKW) [34–36], as follows:σy = ∆σSS + ∆σgb + ∆σOS + ∆σKW . (1)We evaluated the contribution of each strengthening mechanism at peak temperature(800 ◦C for CY1 and 700 ◦C for CY2 and CY3) using Equation (1) above. The contributionof solid solution strengthening, ∆σSS, was calculated using the pySSpredict, a Python-based toolkit that automates the high-throughput calculations of solid solution strength ofcomplex concentrated alloys based on the solid solution strengthening and edge dislocationmodels for FCC and BCC alloys [37,38]. The calculated values of ∆σSS are 20 MPa, 24 MPa,and 32 MPa for the CY1, CY2, and CY3 alloys, respectively. Nb produced the mostprominent strength enhancement via the solid solution strengthening, followed by Tiaddition, which was consistent with the measurement of lattice constants through XRD.Materials 2024, 17, 2280 8 of 12The grain boundary strengthening effect can be expressed as the Hall–Petch relation-ship [39]:∆σgb = σ0 + kd−1, (2)where σ0 is the lattice friction stress, k is the Hall–Petch coefficient, and d is the averagegrain size. Since the average grain size of our alloy is ≈500 µm for CY1, ≈850 µm forCY2, and ≈330 µm for CY3, the effect of grain boundary strengthening is expected to berelatively small in all alloys.In addition to the solid solution strengthening and grain boundary strengthening, theplanar fault energy was assumed to contribute to the significant strength enhancement [40].The strengthening effect of the L12 phase is mainly due to the ordering strengtheningcaused by adding different elements. Since the bonding force between atoms of differentelements is greater than between atoms of the same elements, the ordered arrangement ofatoms of different types will contribute to a higher strength for the ordered alloy, accordingto the expression of the ordering strengthening (∆σOS) [41]:∆σOS = M0.81γAPB2b(3π f8)1/2, (3)where M is the Taylor factor, γAPB is APBE, b is the burgers vector, and f is the volumefraction of the L12 phase. Since APBE dominates the γ′-shearing event by both weakly-coupled dislocations and strongly-coupled dislocations, higher APBE would enhance theoverall strengthening of the L12-ordered phase before the peak temperatures. Chandranand Sondhi [42] investigated the effect of Ti and Nb on the Ni3Al by DFT calculations. Theyreported that Nb and Ti could significantly increase APBE but the Nb effect is stronger withNi3Al1−xNbx with x ≈ 0.20. This research used a DFT method, developed by Cruddenet al. [22], to estimate the APBE of the studied alloys. It is assumed that the change in APBEcan be determined using a linear superposition of the effects of the individual alloyingelements according to the equation∆EAPB = E0APB + ∑ni (kixi), (4)where xi is the concentration in at% of the solute element i in the L12-ordered phase andki is the coefficient for change in APBE, listed in Table 5. E0APB is the APBE for Ni3Al(193 ± 13 mJ m−2) measured using TEM by Kruml et al. [43]. The APBE of the CY1,CY2, and CY3 alloys were 153 mJ/mol, 183 mJ/mol, and 196 mJ/mol, respectively. Thecalculated contribution values to yield strength from ordering strengthening, ∆σOS, are81 MPa, 97 MPa, and 104 MPa for CY1, CY2, and CY3, respectively. The calculation resultindicates that adding Ti or Nb results in an obvious increase in the APBE and, thus, thecontribution of ordering strengthening in the studied alloys.Table 5. Coefficients for change in APBE [22].Coefficient Co Cr Ti Nbki (mJ m−2/at.pct) −1.50 −5.00 15.00 21.40With increasing temperature, the strength of the L12-ordered phase increases, which iscontrolled by the thermally activated cross slip of dislocations from {111} to {010} planes.The number of dislocations in the L12 phase increased as plastic deformation was induced,resulting in the difficulty in dislocation movement along the L12 structure. This studyestimated the contribution of cross-slip-induced strengthening (∆σKW) by subtracting thesolid solution strengthening and the ordering strengthening from the total peak yieldstrength, summarized in Figure 7. The values obtained for cross-slip-induced strengtheningare 325 MPa, 557 MPa, and 554 MPa for CY1, CY2, and CY3, respectively. Therefore, it wasassumed that the cross-slip-induced strengthening dominated the peak yield strength andMaterials 2024, 17, 2280 9 of 12the significant strength enhancement was achieved by adding the fifth elements, namely Tior Nb. Yu et al. [13] have investigated the effect of alloying element on dislocation in Ni3Alusing the DFT method and reported that Ti addition could reduce the cross-slip activationenthalpy and facilitate the cross-slip process to form dislocation locks, thus resulting inmore difficult dislocation movement and higher flow stress in the anomalous temperatureregime of Ni3Al. According to our experimental results, Nb can provide a similar level ofcross-slip-induced strengthening as Ti in the NiCoCrAl alloy system.Materials 2024, 17, x FOR PEER REVIEW 9 of 12   Table 5. Coefficients for change in APBE [22]. Coefficient Co Cr Ti Nb ki (mJ m−2/at.pct) −1.50 −5.00 15.00 21.40 With increasing temperature, the strength of the L12-ordered phase increases, which is controlled by the thermally activated cross slip of dislocations from {111} to {010} planes. The number of dislocations in the L12 phase increased as plastic deformation was induced, resulting in the difficulty in dislocation movement along the L12 structure. This study es-timated the contribution of cross-slip-induced strengthening (𝛥𝜎KW) by subtracting the solid solution strengthening and the ordering strengthening from the total peak yield strength, summarized in Figure 7. The values obtained for cross-slip-induced strengthen-ing are 325 MPa, 557 MPa, and 554 MPa for CY1, CY2, and CY3, respectively. Therefore, it was assumed that the cross-slip-induced strengthening dominated the peak yield strength and the significant strength enhancement was achieved by adding the fifth ele-ments, namely Ti or Nb. Yu et al. [13] have investigated the effect of alloying element on dislocation in Ni3Al using the DFT method and reported that Ti addition could reduce the cross-slip activation enthalpy and facilitate the cross-slip process to form dislocation locks, thus resulting in more difficult dislocation movement and higher flow stress in the anom-alous temperature regime of Ni3Al. According to our experimental results, Nb can provide a similar level of cross-slip-induced strengthening as Ti in the NiCoCrAl alloy system.  Figure 7. Strength contribution of different strengthening mechanisms. In addition to the strength enhancement, a slight decrease in the peak temperature for the positive temperature dependence was observed in the quinary CY2 and CY3 alloys. This phenomenon was also observed in previous studies with ternary alloying addition in Ni3Al alloys and the peak temperature would further decrease with increasing alloying addition of the ternary elements [14,16,17]. Lopez and Hancock [17] suggested that the decrease in peak temperature was due to the strong influence of Ti on the onset of cube slip. Kruml et al. [43] proposed and verified that increasing the CSF energy increased the strength up to the peak temperature and decreased the strength above the peak tempera-ture, thus lowering the peak temperature. According to the CSF energies calculated by Yu et al. [13] using the DFT method, Ti was found to increase the CSF energy, consistent with the lower peak temperature observed in this research. From a technical point of view, the peak temperature is usually at the highest tem-perature where the materials are used [44]. Ni-based superalloys [45–48] and L12-strength-ened HEAs [49–52] have exhibited anomalous positive temperature dependence of Figure 7. Strength contribution of different strengthening mechanisms.In addition to the strength enhancement, a slight decrease in the peak temperature forthe positive temperature dependence was observed in the quinary CY2 and CY3 alloys. Thisphenomenon was also observed in previous studies with ternary alloying addition in Ni3Alalloys and the peak temperature would further decrease with increasing alloying additionof the ternary elements [14,16,17]. Lopez and Hancock [17] suggested that the decrease inpeak temperature was due to the strong influence of Ti on the onset of cube slip. Krumlet al. [43] proposed and verified that increasing the CSF energy increased the strengthup to the peak temperature and decreased the strength above the peak temperature, thuslowering the peak temperature. According to the CSF energies calculated by Yu et al. [13]using the DFT method, Ti was found to increase the CSF energy, consistent with the lowerpeak temperature observed in this research.From a technical point of view, the peak temperature is usually at the highest tempera-ture where the materials are used [44]. Ni-based superalloys [45–48] and L12-strengthenedHEAs [49–52] have exhibited anomalous positive temperature dependence of strengthat intermediate temperatures followed by a decrease in yield strength with increasingtemperatures. The decrease in strength has been explained mainly by degradation ofmicrostructures and transition of deformation mechanisms. Since the L12-ordered singlephase microstructure in our systematically designed alloys was thermodynamically stableat the peak temperatures, the possibility of a microstructure degradation has been ruledout for the cause of the decreased yield strength above the peak temperatures. Our resultsrevealed the intrinsic capability of the multicomponent L12-ordered phases and suggestedthat the difference in the yield strength before the peak temperatures was significantlyinfluenced by the difference in the contribution of cross-slip-induced strengthening, whichwas strongly associated with the alloying compositions of the multicomponent L12-orderedphase. This work has shed light on the significance of cross-slip-induced strengthening inthe multicomponent L12-ordered phase and can serve as a guideline for the future designof L12-strengthened HEAs.According to the DFT calculation result of activation energy of the cross-slip processby Yu et al. [13], Re, W, and Ta possess a higher probability of cross-slip process than thealloying addition of Ti and therefore, these elements are predicted to be more efficient inMaterials 2024, 17, 2280 10 of 12strengthening the Ni3Al phase, which is worthy of further experimental evaluation andinvestigation. In addition, further analysis of the deformed samples will be carried outto elucidate the relationship between the compositions of the L12-ordered phase and theunderlying deformation mechanism.4. ConclusionsThree multicomponent L12-ordered single phase alloys were designed and fabricatedto investigate the effect of multicomponent alloying conditions on the yielding behaviorfrom 500 ◦C to 1000 ◦C. The primary conclusions can be summarized as follows.• Multicomponent L12-ordered alloys, designed using the Calphad Thermo-Calc© soft-ware, were successfully fabricated via VAM and optimized heat treatment. Bothmicrostructure and crystal structure analysis confirmed the formation of L12-orderedsingle phase, demonstrating that ThermoCalc is a powerful software for intermetallicalloy design;• The multicomponent L12-ordered alloys exhibited a positive temperature dependence.The addition of Ti or Nb significantly increased strength up to the peak temperature;• The addition of Ti or Nb enhanced the solid solution strengthening, ordering strength-ening, and cross-slip-induced strengthening. The cross-slip-induced strengtheningwas the most dominant strengthening mechanism in L12-ordered alloys and Ti or Nbaddition remarkably increased the contribution of cross-slip-induced strengthening.Author Contributions: Conceptualization, S.M., H.M., A.-C.Y. and Y.Y.-M.; methodology, H.M.,A.-C.Y. and S.M.; software, C.-Y.W. and S.M.; validation, C.-Y.W.; formal analysis, C.-Y.W., Y.T. andS.M.; investigation, C.-Y.W., Y.T. and S.M.; resources, Y.T. and Y.Y.-M.; data curation, C.-Y.W.; writing—original draft preparation, C.-Y.W.; writing—review and editing, S.M. and Y.Y.-M.; visualization,C.-Y.W.; supervision, H.M., A.-C.Y., S.M. and Y.Y.-M.; project administration, Y.Y.-M.; fundingacquisition, Y.Y.-M. All authors have read and agreed to the published version of the manuscript.Funding: This research received no external funding.Institutional Review Board Statement: Not applicable.Informed Consent Statement: Not applicable.Data Availability Statement: Data are contained within the article.Conflicts of Interest: The authors declare no conflicts of interest.References1. Reed, R.C. The Superalloys: Fundamentals and Applications; Cambridge University Press: Cambridge, UK, 2008.2. Jien-Wei, Y. Recent progress in high entropy alloys. Ann. Chim. Sci. Mat. 2006, 31, 633–648.3. He, J.; Wang, H.; Huang, H.; Xu, X.; Chen, M.; Wu, Y.; Liu, X.; Nieh, T.; An, K.; Lu, Z. A precipitation-hardened high-entropy alloywith outstanding tensile properties. Acta Mater. 2016, 102, 187–196. [CrossRef]4. Shafiee, A.; Moon, J.; Kim, H.S.; Jahazi, M.; Nili-Ahmadabadi, M. Precipitation behaviour and mechanical properties of a newwrought high entropy superalloy. Mater. Sci. Eng. A 2019, 749, 271–280. [CrossRef]5. 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A 2024, 898, 145995.[CrossRef]Disclaimer/Publisher’s Note: The statements, opinions and data contained in all publications are solely those of the individualauthor(s) and contributor(s) and not of MDPI and/or the editor(s). MDPI and/or the editor(s) disclaim responsibility for any injury topeople or property resulting from any ideas, methods, instructions or products referred to in the content.https://doi.org/10.1016/j.matdes.2019.107760https://doi.org/10.1016/S1359-6454(02)00201-Xhttps://doi.org/10.1088/0965-0393/19/2/025008https://doi.org/10.1016/S1359-6454(02)00364-6https://doi.org/10.1016/j.msea.2006.05.032https://doi.org/10.1016/j.jallcom.2021.158878https://doi.org/10.1016/S0921-5093(00)01829-3https://doi.org/10.1016/j.matdes.2017.05.014https://doi.org/10.1016/j.jallcom.2022.165175https://doi.org/10.1007/s11837-015-1484-7https://doi.org/10.1016/j.actamat.2020.02.028https://doi.org/10.1016/j.msea.2022.143712https://doi.org/10.1016/j.msea.2023.145995 Introduction  Materials and Methods  Results and Discussion  Microstructures  Yielding Behavior through Temperatures  Conclusions  References