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[Satoshi Utada](https://orcid.org/0000-0001-6783-9968), Tadaharu Yokokawa, Toshiharu Kobayashi, [Michinari Yuyama](https://orcid.org/0000-0001-7350-2608), Hiroshi Harada, [Toshio Osada](https://orcid.org/0000-0003-1539-9264), [Kyoko Kawagishi](https://orcid.org/0000-0001-7652-9232)

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[Challenging the Limit of Strengthening by γ/γ′ Lattice Misfit on the High Temperature Creep Properties of Ni-Base Single Crystal Superalloy](https://mdr.nims.go.jp/datasets/301df220-6a80-4842-99f3-6380225f0f3c)

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Challenging the Limit of Strengthening by γ/γ′ Lattice Misfit on the High Temperature Creep Properties of Ni-Base Single Crystal SuperalloyTOPICAL COLLECTION: EUROSUPERALLOYS 2026Challenging the Limit of Strengthening by c/c¢ LatticeMisfit on the High Temperature Creep Propertiesof Ni-Base Single Crystal SuperalloySATOSHI UTADA, TADAHARU YOKOKAWA, TOSHIHARU KOBAYASHI,MICHINARI YUYAMA, HIROSHI HARADA, TOSHIO OSADA,and KYOKO KAWAGISHITMS-238MoIr previously achieved a record creep rupture life of 4044 hours at 1100 �C/137 MPa—more than twice that of the sixth-generation Ni-base single crystal superalloyTMS-238—corresponding to a temperature capability of 1136 �C (Larson–Millar conversion by1000 hours rupture life under 137 MPa). This alloy was designed to enhance the high-temper-ature low-stress creep performance of TMS-238 by adding 1.5 at. pct Mo and 2 at. pct Ir. Thisstudy investigates the mechanisms behind this exceptional high-temperature creep performanceand its limitations at other conditions by comparing it to a counterpart alloy, TMS-238MoRu(TMS-238 + 1.5 at. pct Mo and 2 at. pct Ru). At 1100 �C/137 MPa, TMS-238MoIr’s superiorlife was attributed to the combined effects of a refined c/c¢ interfacial dislocation networkachieved through increased magnitude of lattice misfit by Mo addition, and the suppression oftopologically close-packed (TCP) phases by Ir. In contrast, TMS-238MoRu failed in only870 hours due to extensive TCP precipitation. The high c/c¢ coherency stress and decreasedstacking fault energy in both modified alloys induced stacking faults in the c matrix during theinitial heat treatment, before the creep test. These pre-existing faults served as potent nucleationsites for TCP phases at temperatures below 900 �C. Consequently, TCP precipitate-assisteddeformation significantly reduced the creep lives of both concept alloys at 900 �C/392 MPa.These findings demonstrate that lattice misfit tuning must be carefully balanced with stackingfault energy and c/c¢ phase stability to optimize creep performance across a wide range oftemperatures in advanced Ni-base superalloys.https://doi.org/10.1007/s11661-025-08089-5� The Author(s) 2026I. INTRODUCTIONTHE creep properties of a Ni-base single crystal (SC)superalloy for gas turbine application are importantfactors determining component life of turbine bladesand the thermal efficiency of the gas turbine systems.[1–3]To enhance the performance of the aero engines, thetemperature capability of Ni-base SC superalloy hasbeen improved primarily by modification of theirchemical compositions.[1,4] Up to now, from the infor-mation available in literature, third-generation SCsuperalloys containing 4–6 wt pct Re are widely usedin commercial applications.[5–8] Recent trends in thealloy development field aim at more economical alloyswith low or zero Re content, as Re is often considered asexpensive and strategic material.[8–12] These low Realloys attempt to balance the cost and high-temperatureproperties; however, they do not always possesshigh-temperature properties comparable to thehigh-performance alloys such as sixth-generation SCsuperalloy TMS-238 and its advanced version,TMS-238MoIr (TMS-238 + 1.5 at. pct Mo + 2 at. pctIr), developed through the NIMS-ADP (alloy designprogram) approach.[13–17]Despite the general trend against expensive materials,there is still a need to improve the high-temperaturecapabilities of the materials, not only considering thethermal efficiency of the gas turbine systems, but also forspecific operational conditions of aero engines. TheseToshiharu Kobayashi: Deceased author.SATOSHI UTADA, TADAHARU YOKOKAWA, TOSHIHARUKOBAYASHI, MICHINARI YUYAMA, HIROSHI HARADA,TOSHIO OSADA, and KYOKO KAWAGISHI are with theResearch Center for Structural Materials, National Institute forMaterials Science, 1-2-1 Sengen, Tsukuba, Ibaraki, 305-0047, Japan.Contact email: UTADA.Satoshi@nims.go.jpManuscript submitted September 30, 2025; accepted December 11,2025.Article published online January 7, 20261794—VOLUME 57A, MAY 2026 METALLURGICAL AND MATERIALS TRANSACTIONS Ahttp://crossmark.crossref.org/dialog/?doi=10.1007/s11661-025-08089-5&amp;domain=pdfinclude, for example, application of SC superalloys fornon-cooled low-pressure turbine blades,[18] smaller gasturbine for helicopter or private jets that uses very smallblades without cooling systems, and one engine inoper-ative (OEI) events that require contingency poweroperation with overheating.[19–21] Therefore, studies onthe creep properties of SC superalloys at temperaturesabove 1200 �C have become more active.[22–27] More-over, for resource sustainability and stable supply pur-poses, refurbishment and recycling methods for turbineblades have been studied.[28–30] These methods shouldremove barriers for high-end materials to be applied tocommercial components. In fact, a conceptual alloy withcostly Pt addition has been designed and tested aimingat application at such high temperatures.[26,31,32]Ni-base SC superalloys leverage several mechanismsto enhance high-temperature creep properties. The mostfundamental one is to control the fraction of coherent c¢(L12) precipitates to 60–70 pct and their sizes to300-400 nm to hinder dislocation glide.[33] Solid solutionstrengthening of FCC c matrix by adding Re is wellknown as ‘‘rhenium effect,’’ as this element has smallerdiffusion coefficient than other alloying elements, thusslows down dislocation mobility and contributes toextend the creep life.[34,35] However, due to the limitedsolubility of Re in Ni-base superalloys, the precipitationof topologically closed packed (TCP) phases, whichdecrease creep durability of the material, can be an issuewhen the Re is over-alloyed.[36–39] To avoid TCPprecipitation and improve c/c¢ phase stability, Ru, aplatinum group metal (PGM), is added to developfourth and fifth-generation Ni-base SC superalloys withsignificantly improved high-temperature creepstrength.[40–45]Another important parameter determining high-tem-perature low-stress creep durability is the magnitude ofnegative lattice misfit between c matrix and c¢ precipi-tates. Under high-temperature creep loads, cubic c¢precipitates in Ni-base SC superalloys typically trans-form into a plate-like structure known as c¢-raft, whichdevelops transverse to the tensile loading direction.[46–49]When the negative lattice misfit is larger, the c/c¢interfacial dislocation network of the c/c¢-raft becomefiner, which then acts as obstacles for dislocation climband c¢ shearing during the creep deforma-tion.[22,41,46,50–55] By effectively combining c/c¢-misfitstrengthening, solid solution strengthening (by W, Re,and Ta), and enhancing the microstructural stability byRu addition, NIMS developed sixth-generation SCsuperalloy TMS-238 that exhibits excellent creep prop-erties and environmental resistance across wide temper-ature ranges.[13,56] Using the Larson–Miller equation,the temperature capability of TMS-238 is estimated tobe 1117 �C for a 1000-hour creep rupture life at137 MPa.Targeting at improving operating temperature capa-bility of TMS-238 by further advancing c/c¢ misfitstrengthening, TMS-238MoIr was developed by adding1.5 at. pct Mo and 2 at. pct Ir to the base materialTMS-238 in the previous study.[17,57] Mo was added ontop of the original 0.7 at. pct content because it tends topartition into c matrix and increases the magnitude ofnegative lattice misfit[50,58,59]; however, Mo also pro-motes precipitation of TCP phase.[58] Ir addition wasintended to counter this TCP formation and improve c/c¢ phase stability.[17,50,57,60] The developed alloyTMS-238MoIr achieved the longest creep life of4044 hours among all published SC superalloys at1100 �C/137 MPa, a service temperature of 1136 �Cusing the same Larson-Miller conversion for 1000 hourslife under 137 MPa load as summarized in Figure 1.[17]The remaining tasks from the previous studies includea detailed analysis on the mechanism behind theexceptional creep life of TMS-238MoIr under high-tem-perature low-stress creep conditions, and an under-standing of the c/c¢ phase stability and creepperformances under other conditions. To confirm effectof Ir in this alloy system, we also preparedTMS-238MoRu (TMS-238 with 1.5 at. pct Mo and 2at. pct Ru added) as a counterpart. The effects of thesemodifications on creep behaviors, lattice misfit, disloca-tion structure, and TCP phase stability were systemat-ically investigated.II. EXPERIMENTAL DETAILSA. Alloy Design Strategy and Single Crystal SamplePreparationTwo modified TMS-238 variations, TMS-238MoRuand TMS-238MoIr, were prepared as listed in Table I.Al content was adjusted so that the estimated c¢ volumefraction will be 59-61 pct at 1100 �C. The materialsproperties of the experimental alloys, such as c/c¢ latticemisfit (d), solubility index (SI), and density, wereestimated using the NIMS-ADP[14,15,57,61,62] and theyare summarized in Table I. High SI values (SI> 1.2)Fig. 1—Temperature capability of Ni-base superalloys for turbineblade applications, estimated using the Larson-Miller approach for a1000 h creep life at 137 MPa.METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 57A, MAY 2026—1795suggest an increased risk of TCP precipitation duringhigh-temperature exposure,[15,63] although this does notaccount for c/c¢ stabilization effect by Ru or Ir. With a 2at. pct Ru addition, this SI limit increases from 1.2 to1.3, meaning that the composition of the baselineTMS-238 is within an acceptable range.[63] Accordingto the estimation, additional Mo can be very efficient inincreasing the magnitude of negative c/c¢ lattice misfit.However, simultaneously, increased SI values give awarning that there is very high chance of early TCPnucleation.The single crystal materials were fabricated fromhigh-purity raw materials using the withdrawing preci-sion casting method in a directional solidificationvacuum induction melting furnace with a withdrawalrate of 200 mm/h. The cast bars have a 10.5 mmdiameter and a 130 mm length. The materials underwentsolution heat treatment and subsequent two-step agingtreatments using a vacuum furnace with the conditionspresented in Table II. For all heat treatments, materialswere moved to a cooling chamber and cooled down bythermal radiation. The cooling rate from solutiontemperature was about 300 K/min. The aging condi-tions were carefully selected to yield similar average c¢precipitate sizes across the three alloys. Edge lengthsalong [100] direction of cuboidal c¢ precipitates weremeasured for at least 150 precipitates to determine theaverage precipitate size of each material.B. High-Temperature TestsThe crystallographic orientation of heat-treated SCbars was measured by the back-reflection Laue method,and the bars with less than 5 deg of longitudinalcrystallographic misorientation from< 001>wereselected for a tensile creep test performed in this study.The creep testing specimens have a 4 mm diameter and a20 mm parallel section, in accordance with theJIS-Z2271 standard.[64] Creep tests were performedfollowing the same standard at 1100 �C/137 MPa,1000 �C/245 MPa, 900 �C/392 MPa, and 800 �C/735 MPa for all three alloys listed in Table I.Phase stability and equilibria analyses were performedon the SC billets using the same prolonged aging testmethod from previous studies.[60,64] The SC billets witha 5 mm thickness were cut from as-heat-treated SC bars.To obtain a time-temperature-transformation (TTT)diagram, the cut-out billets were further exposed in aresistive furnace at 800-1200 �C for up to 1000 hours.To analyze phase equilibria at the creep test tempera-ture, severe plastic deformation was introduced into abillet at room-temperature, followed by annealing at1100 �C for 500 hours. This strain aging method wasapplied to the material to enhance recrystallization andfaster microstructure evolution, enabling the c/c¢microstructure to approach equilibrium and becomecoarse enough to measure their individual compositionusing microscopy.[65]C. Microstructure ObservationsMicrostructures of heat-treated, creep-ruptured, andprolong-aged samples were observed using a scanningelectron microscopy (SEM). A field emission gun-scan-ning electron microscope (FEG-SEM), Zeiss Gemini300, was operated for the observations. Superalloysamples for microstructure observation were mountedinto a conductive resin, followed by metallographicpolishing to colloidal silica with a 0.05 lm particle size.Because each specimen condition required examinationof different microstructural features, the observationplane was varied accordingly. As-heat-treated sampleswere polished on planes near (001) plane, creep-rupturedspecimens were sectioned and polished in the loadingdirection (crystallographic plane unknown), and pro-longed-aged samples were polished on planes parallel tothe growth direction.Creep-ruptured samples were analyzed using electronbackscatter electron diffraction (EBSD) by EDAXCMOS camera equipped on the Zeiss Gemini 300FEG-SEM, with a step of 50 nm. Indexed data werepost-processed using EDAX OIM Analysis� 8. Thesame FEG-SEM with angular selective backscatter(AsB) detector was used to obtain electron channelingcontrast images (ECCIs) operated at an accelerationvoltage of 20 kV and probe size of 60 lm. LocalTable II. Heat Treatment Conditions Applied to theExperimental MaterialsSolution 1st Aging 2nd AgingTMS-238 1345 �C, 20 h 1150 �C, 2 h 870 �C, 20 hTMS-238MoRu 1340 �C, 20 h 1100 �C, 4 h 870 �C, 20 hTMS-238MoIr 1340 �C, 20 h 1150 �C, 4 h 870 �C, 20 hTable I. Chemical Composition and Estimated Properties of the Experimental AlloysComposition (Ni-Bal) c/c¢ Lattice Misfitat 1100 �C (Pct)SI Density(g/cm3)Co Cr Mo W Al Ta Hf Re Ir RuTMS-238 at. pct 7.0 5.6 0.7 1.4 13.9 2.7 0.04 2.2 3.1 � 0.34 1.25 8.99wt pct 6.5 4.6 1.1 4.0 5.9 7.6 0.1 6.4 5.0TMS-238MoRu at. pct 7.0 5.6 2.2 1.4 13.1 2.7 0.04 2.2 5.1 � 0.50 1.50 9.09wt pct 6.4 4.5 3.2 3.9 5.45 7.4 0.1 6.2 8.0TMS-238MoIr at. pct 7.0 5.6 2.2 1.4 13.1 2.7 0.04 2.2 2.0 3.1 � 0.41 1.48 9.36wt pct 6.2 4.4 3.1 3.8 5.3 7.2 0.1 6.1 5.8 4.81796—VOLUME 57A, MAY 2026 METALLURGICAL AND MATERIALS TRANSACTIONS Acompositions of c and c¢ phases were quantitativelymeasured using an Electron-probe Microanalyzer(EPMA) Shimadzu EPMA-1610 at the recrystallizedand coarsened microstructure.To analyze the effect of c/c¢ interfacial misfit,creep-ruptured samples were observed using a transmis-sion electron microscopy (TEM), JEOL JEM-4010 andJEM-2100. End of the parallel gauge section close totransition radius were cut into discs perpendicular to theloading direction, followed by mechanical polishing forthinning down to< 100 lm. The thinned discs wereelectrochemically polished by twin jet method (Fis-chione Instruments) with following conditions: A 40 Velectrolytic voltage on a solution ofCH3COOH:HCO4 = 10:1 at a 10 �C solutiontemperature.III. RESULTSA. Microstructure Before the High-Temperature TestsMicrostructure after the solution and aging heattreatments was observed using SEM in back scatterelectron (BSE) imaging mode as shown in Figure 2. Allthe experimental alloys have been successfully solutiontreated and present a microstructure without c/c¢ eutec-tic pool. All the materials present cuboidal c¢ precipi-tates in general, which have an edge length average of211.3 nm (TMS-238), 199.5 nm (TMS-238MoRu), and241.6 nm (TMS-238MoIr) as presented in Figure 3.Unlike standard TMS-238, the conceptualTMS-238MoRu and TMS-238MoIr exhibit thick cchannels (~ 50 nm) and fine c¢ precipitates (~ 25 nm)inside these channels. Additionally, in these modifiedalloys, some precipitates appear to have merged withnearby ones through very thin channels. While theaverage precipitate sizes are similar across all alloys,these characteristics of precipitates in the conceptualalloys are reflected in error bars in Figure 3. In additionto the c/c¢ phases, TMS-238MoRu and TMS-238MoIrhave diagonal traces as shown in Figures 2(b) and (c),respectively, and are surrounded by dotted curves. Theseareas represent thick c channels lying on the observationplane, and precipitates beneath this c-channel layer areexposed due to relatively high acceleration voltage(20 kV).These c-channel layers were observed at highermagnification using ECCI and the obtained imagesshown in Figure 4. Within the thick c-channels, asubstantial number of dislocations and planar faultswere observed for both TMS-238MoRu andTMS-238MoIr. Not only were they observed in thematrix, but also some planar faults in the matrix wereextending into c¢ precipitates (examples are pointed byarrows in Figure 4(d)). These are mainly stacking faultsas they lie on the {111} plane and continuouslyextending from the matrix.[66–68]B. Creep Properties of TMS-238 and Concept VariantsThe results of creep tests on the three experimentalalloys, TMS-238, TMS-238MoRu, and TMS-238MoIr,at 1100 �C/137 MPa, 1000 �C/245 MPa, 900 �C/Fig. 2—SEM-BSE microstructures of TMS-238 (a), TMS-238MoRu (b), and TMS-238MoIr (c) before high-temperature tests. The dotted curveoutlines an example of thick c-channel oriented parallel to the imaging plane close to (001) plane.Fig. 3—Edge length of c¢ precipitate in the dendritic core region ofTMS-238, TMS-238MoRu, and TMS-238MoIr before thehigh-temperature tests.METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 57A, MAY 2026—1797392 MPa, and 800 �C/735 MPa are shown in Figure 5along with two representative commercial alloys,CMSX-4 and CMSX-10.[5,69] The creep strain rate isplotted against the creep time as shown in Figure 6.At 1100 �C/137 MPa, compared with baselineTMS-238 (1284 and 1925 hours), TMS-238MoIrdemonstrated superior creep durability that can reachup to 4044 hours.[17] As mentioned earlier, this is thelongest creep life reported for this condition in theavailable literature.[17] TMS-238MoIr has very similarcreep properties to TMS-238 at 1000 �C/245 MPa;however, this material has about half the life of thebaseline at the lower temperature conditions of 900 �C/392 MPa and 800 �C/735 MPa. TMS-238MoRu exhib-ited the lowest creep life of these three materials in allfour conditions tested, underperforming in comparisonwith the baseline TMS-238.Creep strain rate curves in Figure 6 explain theirdifferences in detail. At 1100 �C/137 MPa and 1000 �C/245 MPa (Figures 6(a), (b)), TMS-238MoIr andTMS-238 have very similar trend during secondarycreep stage, whereas TMS-238MoRu was deforming athigher rate in this stage. At 900 �C/392 MPa and800 �C/735 MPa (Figures 6(c), (d)), TMS-238MoRushowed immediate transition into the creep acceleratingstage in comparison to the other two alloys.Larson-Miller parameters of TMS-238,TMS-238MoRu, and TMS-238MoIr calculated fromtheir creep rupture life are plotted in Figure 7. Com-pared to the commercial alloys, CMSX-4 andCMSX-10, TMS-238 has well balanced creep lives forall temperature ranges. Increasing magnitude of nega-tive lattice misfit with suppression of TCP phase,approach tested in this study, has shown effectivenessat creep temperatures above 1000 �C.C. Microstructural Analyses After Creep Testsat 1100 �C/137 MPaPost-mortem microstructures of creep-ruptured sam-ples were analyzed to investigate the mechanism of theimproved and degraded creep properties ofTMS-238MoIr and TMS-238MoRu compared toTMS-238 at the targeted condition of at 1100 �C/Fig. 4—Detailed microstructures of TMS-238MoRu (a), (b) and TMS-238MoIr (c), (d) obtained by ECCI. The dotted curve outlines an exampleof thick c channel oriented parallel to the imaging plane and the dotted lines are direction of planar faults inside the thick c channel. Arrows arepointing at examples of stacking faults in c¢ which are extending from the c channel.1798—VOLUME 57A, MAY 2026 METALLURGICAL AND MATERIALS TRANSACTIONS A137 MPa. Figure 8 presents the microstructures of eachalloy after the creep rupture test at 1100 �C/137 MPa.Figures 8(a) through (c) show the microstructure at thedendritic core located 1 mm away from the fracturesurface of TMS-238 (ruptured at 1284 hours),TMS-238MoRu (ruptured at 870 hours), andTMS-238MoIr (ruptured at 4044 hours) samples,respectively. Representative c¢-rafted microstructure ofthe same samples at the dendritic core close to the gageend (0.5 mm away from the end toward fracture surface)are shown in Figures 8(d) through (f).The baseline TMS-238 has some TCP precipitates(approximately 0.3 pct by area) that form surroundingc¢-envelope, which can be a path for dislocation glideand climb during the creep deformation. Figure 8(a) is aparticular example of creep void growing at c¢-TCPregion which is speculated to be accelerated by thec¢-envelope. In contrast, TMS-238MoRu (Figure 8(b))exhibited a substantial precipitation of TCP phase (areafraction ~ 6.5 pct) and coarsening of c¢-rafts aroundTCP precipitates. Surprisingly, even after 4000 hours ofexposure, only few TCP precipitates were nucleated inTMS-238MoIr that has negligible effect on the creep lifeand failure processes. The failure mechanism ofTMS-238MoIr seemed to be typical nucleation andgrowth of creep voids (see example in Figure 8(c)).Superior c/c¢ phase stability during the creep testing hasbeen achieved by substitution of Ru by Ir.c¢-rafted structure shown in Figures 8(d) through (f)are typically observed in Ni-base SC superalloys creeptested at temperatures above 900 �C. Acquisition loca-tions for these images were 0.5 mm from the gage end toavoid effect of severe straining near the fracture surfacethat can degrade c¢-rafts. In all experimental alloys, acontinuous raft structure oriented perpendicular to thestress axis was observed. To compare the c¢-rafts and c¢volume fraction of these materials, SEM-BSE imageswere analyzed using ImageJ Fiji[70] with Trainable WekaSegmentation[71] and LinearDistance[72] plugins, and theresults are shown in Figure 9. The difference of theFig. 5—Creep curves of TMS-238, TMS-238MoRu, and TMS-238MoIr conducted at 1100 �C/137 MPa (a), 1000 �C/245 MPa (b), 900 �C/392 MPa (c), and 800 �C/735 MPa (d). Two representative commercials alloys, CMSX-4 and CMSX-10, are included for comparison.METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 57A, MAY 2026—1799c¢-rafts between these materials are within the statisticalscatter and not significant overall. Volume fraction of c¢phase was higher at the interdendritic area for all thetested material, due to the dendritic segregation. Thissuggests that the magnitude of segregation was differentamong the materials; particularly, TMS-238MoIrshowed the largest deviation between the primarydendrite core and the interdendritic area, probably byeffect of slower diffusion of Ir. Average volume fractionwas 64 to 65 pct, which is slightly higher than theestimation by NIMS-ADP (Table I), but consistent forall the materials as expected.To validate the concept of strengthening via very highnegative c/c¢ lattice misfit, c/c¢ interfacial dislocationnetworks on the plane perpendicular to the loadingdirection were examined using TEM, and representativemicrographs are shown in Figure 10. Dislocationspacing was quantified using the line-intercept methodby counting the intersections between dislocations andlines drawn along the [100] direction, allowing directcomparison of dislocation density.[64,73] A minimum of70 dislocation segments were measured for statisticalreliability, and the results are summarized in Figure 11.The data clearly show a refinement of the interfacialdislocation network in the following order: TMS-238(38.2 nm), TMS-238MoIr (25.5 nm), andTMS-238MoRu (19.7 nm). These results demonstratethat increasing the magnitude of negative lattice misfitthrough alloy composition modification (addition of 1.5at. pct Mo) effectively promotes dislocation networkFig. 6—Creep strain rate curves of TMS-238, TMS-238MoRu, and TMS-238MoIr conducted at 1100 �C/137 MPa (a), 1000 �C/245 MPa (b),900 �C/392 MPa (c), and 800 �C/735 MPa (d).Fig. 7—Larson–Miller plots (C = 20) for Ni-base SC superalloyscreep tested in this study.1800—VOLUME 57A, MAY 2026 METALLURGICAL AND MATERIALS TRANSACTIONS Arefinement. With the combination of c/c¢ microstructurestability, the longest creep life was achieved byTMS-238MoIr.D. Microstructural Analyses After Creep Testsat 900 �C/392 MPaContrarily to the successful creep property improve-ment at 1100 �C/137 MPa by tailoring c/c¢ lattice misfit,creep life of two modified alloys drastically decreasedfrom that of the baseline TMS-238 at 900 �C/392 MPacondition. Microstructures were analyzed using SEMand EBSD to investigate the cause of the creep lifereduction.Color coded inverse pole figure (IPF) maps and grainreference orientation deviation (GROD) maps wereobtained from samples after creep rupture tests at900 �C/392 MPa (Figure 12). The EBSD data wasFig. 8—SEM-BSE images of TMS-238 (a), (d), TMS-238MoRu (b), (e), and TMS-238MoIr (c), (f) samples after creep rupture test at 1100 �C/137 MPa. Observed locations are 1 mm from the fracture surface toward gage end (a) to (c) and 0.5 mm from gage end toward fracture surface(d) to (f).Fig. 9—(a) c¢-raft width (perpendicular to the loading axis) and thickness (parallel to the loading axis) and (b) c¢ volume fraction at primarydendritic core and interdendritic area of TMS-238, TMS-238MoRu, and TMS-238MoIr samples after creep rupture tests at 1100 �C/137 MPa.Error bars represent standard deviation.METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 57A, MAY 2026—1801collected around creep void locating 200 lm away fromthe fracture surface. Notable characteristics of themodified alloys are that TCP precipitates are spreadingwidely in the microstructure, and chains of recrystallizedgrains are extending from the edges of the creep voids.In those recrystallized chains, TCP phases are precipi-tating in granular form (see Figure 13). Isolated voidnucleation and crack propagation are observed alongthe chain of recrystallized grains (see Figure 12(c)).These creep voids and recrystallized chains wereobserved up to ~ 2 mm away from fracture surfaceand not at the regions near the gage end. Therefore, voidnucleation and recrystallization occurred during thecreep acceleration stage of TMS-238MoRu andTMS-238MoIr, and the failure process was by far fasterand shorter than that of the baseline TMS-238 (see alsocreep strain rate in Figure 6(c)).GROD maps in Figure 12 represent the local crystal-lographic misorientation within individual grains, cal-culated with respect to the orientation of a pixel withminimum kernel average misorientation of each grain.GROD quantifies the degree of intragranular latticedistortion, which means the higher GROD misorienta-tion values typically indicate regions of increased plasticdeformation. Creep void in TMS-238 is obviously thecrack initiation site and higher magnitude of latticerotation around the void induced twinning. In the alloyswith significant TCP precipitation (TMS-238MoRu inFigure 12(d) and TMS-238MoIr in Figure 12(f)), theGROD maps reveal the plastic deformation accumula-tion at TCP interfaces. This localized deformation likelycauses the higher creep strain rates observed inTMS-238MoRu throughout the entire test and earlieronset of tertiary creep in TMS-238MoIr (Figure 6(c)).The local crystallographic misorientation is recoveredby the formation of recrystallization chain.Regions far from the fracture surface were observedand shown in Figure 14. TMS-238MoRu, shown inFigure 14(b), exhibited a fine but higher number of TCPprecipitates compared to the other two alloys, whichcorrelates with its significantly reduced creep life. In thehigher magnification of microstructure ofTMS-238MoIr (Figure 14(d)), TCP phase is precipitat-ing from c-channel and defects inside the c¢ (probablystacking faults) are extending from those TCP precip-itates. Shape of c-channels connected to TCP precipitateare distorted along the growth direction of TCP.E. c/c¢ Phase Stability and TCP Phase PrecipitationAs the creep properties of the experimental alloyswere affected by TCP phase precipitation, the differencesin c–c¢ microstructural stabilization effects resultingfrom Mo, Ru, and Ir additions in Ni-base superalloyTMS-238 need clarification. The c–c¢ partitioningbehavior of alloying elements in TMS-238MoRu andTMS-238MoIr was investigated from an equilibriumperspective, while the isothermal transformation associ-ated with TCP precipitation was examined from akinetic perspective.The partition ratio K of alloying elements was definedas K = Xc/Xc¢, where Xc and Xc¢ represent the chem-ical composition in the c phase and c¢ phase, respec-tively. Figure 15 shows the microstructure ofTMS-238MoRu and TMS-238MoIr after the strainaging at 1100 �C for 500 hours, having the TCP phaseFig. 10—TEM micrograph of TMS-238 (a), TMS-238MoRu (b), and TMS-238MoIr (c) samples after creep rupture test at 1100 �C/137 MPaobtained with beam direction parallel to [001] zone axis, showing c/c¢ interfacial dislocation networks.Fig. 11—Average c/c¢ interfacial dislocation network spacing along[100] direction of TMS-238, TMS-238MoRu, and TMS-238MoIrsamples after creep rupture tests at 1100 �C/137 MPa. Error barspresent standard deviation.1802—VOLUME 57A, MAY 2026 METALLURGICAL AND MATERIALS TRANSACTIONS AMETALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 57A, MAY 2026—1803precipitated in both alloys under nearly equilibriumstate. Figure 16 presents the partition ratios of thealloying elements between c and c¢, revealing marginaldifferences in partitioning behavior betweenTMS-238MoRu and TMS-238MoIr compared toTMS-238. This indicates that, thermodynamically, par-titioning of alloying elements does not vary by additionof Mo and Ru or Ir to TMS-238. Overall partitioningtrends are similar to those have been reported in theprevious studies.[65,74,75]TTT diagrams for TMS-238MoRu andTMS-238MoIr are presented in Figure 17. In thesediagrams, circles and triangles indicate the samplingpoints for TCP phase precipitation in TMS-238MoRuand TMS-238MoIr, respectively, with filled symbolsrepresenting conditions under which TCP phases wereobserved. The solid curve was drawn based on con-firmed TCP precipitation points. For comparison, theTTT diagram for TMS-238 was drawn using data fromthe previous study.[60] The nose temperature for bothTMS-238MoRu and TMS-238MoIr is in between900 �C and 1000 �C, and the time of first TCP precip-itation shifts to shorter aging durations for theTMS-238MoRu. This trend is consistent regardless ofobservation locations; dendritic core (Figure 17(a)) andinterdendritic (Figure 17(b)) regions. In the dendriticcore region, TCP phase precipitation in TMS-238MoRuat 900–1000 �C occurred earlier than 50 hours of aging.From both equilibria and kinetic perspectives, theaddition of Mo to TMS-238 promotes TCP phaseformation, even in the presence of TCP-suppressingelements such as Ru and Ir. In addition, although bothTMS-238MoRu and TMS-238MoIr have higher ten-dency for early TCP precipitation compared toTMS-238, replacing 2 at. pct of Ru by Ir effectivelydelays TCP formation, thus stabilizing the c/c¢microstructure, particularly at temperatures higher than1050 �C.IV. DISCUSSIONThis discussion focuses on how targeted alloyingmodifications to the sixth-generation Ni-based singlecrystal superalloy TMS-238 influence the c/c¢ latticemisfit, and how these changes affect microstructuralevolution and creep deformation behaviors at 1100 �Cand 900 �C. Particular emphasis is placed on the role ofc/c¢ misfit in controlling interfacial dislocation networksand phase stability. The influence of Ir addition—espe-cially in comparison to Ru—is further discussed in thecontext of its impact on high-temperature properties.These findings are compared to previous studies con-ducted at NIMS, providing a comprehensive under-standing of how elemental partitioning, diffusionbehavior, and phase stability interact to control creepperformance of Ni-base SC superalloys.A. Effect of Alloy Composition Modifications on the c/c¢Lattice Misfit and the Creep Properties at 1100 �C/137 MPaTailoring the c/c¢ lattice misfit and to densify theinterfacial dislocation network is a key strategy forenhancing the high-temperature creep resistance ofNi-base SC superalloys. The observed dislocation net-work after the creep test showed very fine structure asquantified in Figure 11. This confirms that the conceptalloys with addition of 1.5 at. pct Mo have increased themagnitude of negative lattice misfit, particularly effec-tive in the case for TMS-238MoRu with the smallestdislocation spacings. The elemental partitioning side ofMo in both TMS-238MoRu and TMS-238MoIr did notchange by addition of Ru or Ir (see Figure 16).TMS-238MoRu had the highest partition coefficient ofMo to c among three alloys, and Ru also preferentiallypartitions into c side; both can contribute to increasinglattice constant of c phase.[50,58,59,76–78] The Ir partition-ing to c¢ observed in TMS-238MoIr, similar trendreported in the literature, possibly increased the latticeconstant of c¢, leading to slightly coarser interfacialdislocation network of this alloy.[74–77,79]The primary creep stage of the creep curves from thecreep tests at 1100 �C/137 MPa is shown in Figure 18.The time required to reach the steady-state creep regionfrom the onset of creep was shortest forTMS-238MoRu, followed by TMS-238MoIr andTMS-238. These results are consistent with previousreports explaining the relationship that a more negativelattice misfit accelerates c/c¢ phase rafting, and therefore,bFig. 12—IPF color-coded orientation maps referring to the tensiledirection (a), (c), (e) and GROD maps respect to the orientation ofpixel with minimum kernel average misorientation (b), (d), (f) ofTMS-238 (a), (b), TMS-238MoRu (c), (d), and TMS-238MoIr (e), (f)after creep rupture tests at 900 �C/392 MPa. Data were collected atthe location 200 lm away from fracture surface. (Color imageavailable online).Fig. 13—SEM-BSE image of TMS-238MoRu after creep rupture testat 900 �C/392 MPa (same location as Fig. 11(c), (d)). Arrows arepointing at examples of TCP phase precipitates in the chain ofrecrystallized grains.1804—VOLUME 57A, MAY 2026 METALLURGICAL AND MATERIALS TRANSACTIONS Athe primary creep stage finishes earlier. Normally in theNi-base SC superalloys, morphology of c¢-raft alsodepends on the c/c¢ lattice misfit and initial c¢ precipitatesize, which influence the microstructure evolution duringthe primary creep stage at temperatures above1000 �C.[46,80,81] The perfection degree (aspect ratio) ofc¢-raft is important for creep life since the smoother andlonger c¢-raft means wider net area of horizontal c/c¢interface that limits dislocation path.[46] However, thequantified width and thickness of c¢-raft were almost thesame for all the tested alloys (see Figure 9), andtherefore the perfection degree of the c¢-raft did notcontribute to the varying creep life. It is worth notingthat c¢-rafts observed in this study are by far thinnerthan those in other commercial superalloys reportedpreviously. This suggests that the dislocation networkdensity is one of dominant factors determining creepdurability of these alloys. The exceptional creep life of4044 hours of TMS-238MoIr at 1100 �C/137 MPa cantherefore be attributed to the synergistic effect of a highc/c¢ lattice misfit and improved c/c¢ phase stability thatprevented TCP phase precipitation.Extremely high magnitude of c/c¢ lattice misfit alsoinfluenced the microstructure before the creep tests.With 0.3 pct misfit (normally seen for commercial alloyswith 3–5 wt pct Re), the expected c/c¢ coherency stressarising from the lattice misfit can be up to350 MPa.[82,83] TMS-238MoRu has dislocation spacingof ~ 20 nm, which is nearly half of TMS-238 (seeFigure 11). The relationship between absolute value oflattice misfit d, dislocation network spacing d, andmagnitude of Burgers vector b can be simplified in thefollowing Eq. [1][78,84]:dj j ¼ bj j=d ½1�Fig. 14—SEM-BSE images of TMS-238 (a), TMS-238MoRu (b), and TMS-238MoIr (c) after creep rupture test at 900 �C/392 MPa. Images weretaken at dendritic core region located 0.5 mm from gage end toward fracture surface. (d) Is detailed image of (c) obtained by ECCI and arrowsin this image indicate defects extending from TCP precipitates.METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 57A, MAY 2026—1805With a rough assumption that magnitude of Burgersvector for edge component is similar for TMS-238 andTMS-238MoRu, the c/c¢ lattice misfit ofTMS-238MoRu is estimated to be nearly twice ofTMS-238 because the dislocation spacing is about thehalf (Note that the misfit values in Table I areestimations from NIMS-ADP and are not derived froman experimental result). Therefore, much highercoherency stress at the c/c¢ interface is expected for thealloys with Mo addition. At the aging temperature of870 �C, volume fraction of c¢ is nearly maximum leveland the compressive stress in the c matrix is expected toincrease drastically.Moreover, increasing Mo and Ru content naturallydecreases stacking fault energy of c phase.[85–87] Bothcoherency stress and decreased stacking fault energyFig. 15—SEM-BSE images of TMS-238MoRu (a) and TMS-238MoIr (b) after strain aging at 1100 �C—500 h, showing c/c¢ and TCP phases.Fig. 16—Elemental partitioning ratios for alloying elements favoring partition into the c phase (a) and the c¢ phase (b), based on compositionalanalysis of samples strain-aged at 1100 �C for 500 h followed by EPMA. Partitioning ratios of Ru and Ir are included in (b).1806—VOLUME 57A, MAY 2026 METALLURGICAL AND MATERIALS TRANSACTIONS Acontributed to the formation of stacking faults inside thematrix observed by ECCI in Figure 4. Matrix channelshave variation of thicknesses; however, the stackingfaults extending from very thin channel (as thin as10 nm) can be observed in the same Figure 4. AlthoughMo and Ru are known to decrease the stacking faultenergy, the formation of stacking faults in the as-heat-treated condition (before mechanical loading) has notbeen previously reported. This suggests that coherencystress plays a significant role in the generation of thesestacking faults.While pre-existing faults may accelerate the transitionfrom primary to secondary creep stage, they should notaffect the creep deformation behaviors when the c¢-raft-ing occurs rapidly (such as> 1100 �C). This is becausematrix plasticity is a natural process that first occursduring the primary creep stage. Hence, the relaxation ofthese planar faults or matrix dislocations into theinterfacial dislocation network will progressanyway.[88–90]B. Influence of Ru and Ir on the c/c¢ Phase Stability,TCP Formation, and Creep Properties of Ni-Base SCSuperalloyAmong PGMs, Ir has an ability to suppress theprecipitation of TCP phases when substituted for Ru inNi-base SC superalloys.[50,65,91] Notably, Ir has a supe-rior TCP suppression effect in TMS-238 compared toRu.[60,79] Based on results in the present study, thesuppression of TCP phase precipitation inTMS-238MoIr is attributable not to a change ofthermodynamic equilibrium such as improvement ofsolubility limit in Ni-base superalloys by Ir addition, butrather reducing elemental mobility by very low diffusiv-ity of Ir in Ni alloys than other elements.[92,93] This isalso consistent in the case where Ru of TMS-238 wassimply replaced by Ir.[60,74] Thus, the superior creeprupture life of TMS-238MoIr at 1100 �C/137 MPaderives from the combined effects of increased negativelattice misfit by Mo addition that leads to the formationof a refined c–c¢ interfacial dislocation network inducedand enhanced microstructural stability provided by Iraddition. Slower diffusion of Ir itself among alloyingelements can also contribute to improving the creeplife.[92,93] In contrast, formation of long and extendingTCP precipitates during the creep of TMS-238MoRu(see Figure 8(b)) likely accounts for its reduced creeprupture life at the same condition 1100 �C/137 MPa.At 1000 �C/245 MPa, the benefits of the Ir additionwere less pronounced, with TMS-238MoIr exhibitingcreep life merely comparable to the baseline TMS-238.This intermediate performance can be due to thedendritic segregation dependent phase stability. WhileFig. 17—Time–temperature–transformation (TTT) diagrams for TCP phase precipitation in TMS-238MoRu and TMS-238MoIr, microstructureobserved at dendritic core (a) and interdendritic (b) regions. The TTT curve for TMS-238 is included for reference, based on data from Ref.[60]Circles and triangles indicate sampling points for TMS-238MoRu and TMS-238MoIr, respectively; filled symbols represent conditions whereTCP phases were observed.Fig. 18—Initial creep strain evolution of TMS-238, TMS-238MoRu,and TMS-238MoIr alloys tested at 1100 �C/137 MPa, highlightingthe transition from primary to steady-state creep behavior.METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 57A, MAY 2026—1807TCP precipitation was widespread in TMS-238MoRu, itwas confined to the dendritic core regions inTMS-238MoIr (microstructural analyses of creep-rup-tured samples, not shown here). This TCP precipitatedistribution is consistent with the TTT diagrams(Figure 17). This suggests that while the kinetic sup-pression by Ir is effective in the interdendritic regions atthis temperature, it is insufficient to prevent TCPnucleation in the dendritic cores where the segregationof Re is stronger.Previous studies have shown that the slower diffusionof Ir and its ability to delay TCP phase formation cancontribute to improving the creep rupture life of Ni-baseSC superalloys at intermediate temperature conditions,such as 900 �C/392 MPa.[74,79] However, this was notthe case in the present study due to the severe TCP phaseprecipitation, as observed in Figures 13 and 14. Theprimary reason for this is the lack of matrix solubilitycaused by the high Mo content, as suggested by high SIvalue listed in Table I. Another likely reason is the initialmicrostructure, which contained a high density ofstacking faults (Figure 4). These faults lie on {111}planes, which are also the preferred crystallographicplanes for TCP precipitate growth. The post-mortemmicrostructure reveals that TCP phases and c¢-shearingfaults are on the same crystallographic planes (seeFigure 14(d)). Considering that atom probe studies thathave shown the segregation of solutes such as Re and Wat faults,[94–96] it is plausible that these stacking faultsand associated planar faults in c¢ precipitates shouldfavor similar segregation, and consequently facilitatesTCP nucleation and growth from stacking faults in cmatrix.[97] Therefore, extremely high lattice misfit anddecreased stacking fault energy adversely affected thecreep properties of the material. This detrimentalmechanism appears to be activated at the creep temper-atures below 900 �C (including creep test at 800 �C/735 MPa), where the temperature is too low for signif-icant relaxation via c¢-rafting. Furthermore, under thesecreep conditions, the high applied stress promotesextension of pre-existing planar faults into the c¢ phase,accelerating the onset of tertiary creep and leading topremature failure.C. Ni-Base SC Superalloy Design Including IrFigure 19 summarizes the relationship between themisfit and creep rupture life at 1100 �C and 137 MPa forfirst to sixth-generation Ni-base SC superalloys. Themisfit of TMS-238MoIr was estimated by using the sameEq. (1). The circles represent alloys developed in NIMSand the triangles represent commercial alloys. The trendof the magnitude of the negative lattice misfit and creeprupture life under high-temperature and low-stressconditions are consistent for the conceptualTMS-238MoIr alloy presented here. TMS-238MoRuwith even higher magnitude of the misfit does not followthe trend. This is because of unstable microstructurewith TCP precipitates. This observed trend—achieving ahigher lattice misfit requires Mo addition, which in turnincreases the risk of TCP formation—is consistent withthe findings of Bezold et al.[88]Unbalanced creep properties are because of high Mocontent for improving lattice misfit and decreasingstacking fault energy, which induces high risk of TCP,and hence the creep performance at lower temperaturecannot be sustained. Ir on the other hand, has shownability not only to prevent TCP formation, but also toimprove creep performance at intermediate to lowertemperatures by decreasing diffusivity and increasingsolid solution effect.[74,79,98] Further studies are neces-sary to optimize the materials’ creep properties acrosstemperature ranges.It needs to be mentioned that critical problems of thismaterial are material cost (150 USD/g[99]) and density(approximately 9.37 g/cm3 for TMS-238MoIr by ADP,see Table I). However, studies on Ni-base superalloyswith Ir for both blade alloy and coating applicationshave shown very small addition of Ir improves oxidationand corrosion resistances.[100,101] Ir indeed has positiveeffect on the high-temperature properties of Ni-base SCsuperalloys in overall. Therefore, further optimizationmay open a door for future application of this class ofmaterial.V. CONCLUSIONTargeting the improvement of high-temperaturelow-stress creep properties of sixth-generation Ni-basesingle crystal superalloy TMS-238, TMS-238MoIr(TMS-238 + 1.5 at. pct Mo + 2 at. pct Ir) was previ-ously developed and demonstrated the world’s longestcreep life of 4044 hours at 1100 �C/137 MPa. This studyprepared counterpart alloy TMS-238MoRu(TMS-238 + 1.5 at. pct Mo + 2 at. pct Ru) to inves-tigate creep properties and c/c¢ phase stability of thethree alloys including baseline TMS-238. The followingmain conclusions can be established:Fig. 19—Relationship between c/c¢ lattice misfit and creep rupturelife at 1100 �C/137 MPa. TMS-238MoIr exhibits an excellent creeprupture life beyond TMS-238.1808—VOLUME 57A, MAY 2026 METALLURGICAL AND MATERIALS TRANSACTIONS A� In contrast to TMS-238MoIr, which exhibited thehighest temperature capability of 1136 �C, which isbeyond that of TMS-238 at 1120 �C (Larson–Millarconversion by 1000 hours rupture life under137 MPa), TMS-238MoRu underwent extensiveTCP phase precipitation and showed a significantlyreduced creep life of 870 hours—about half that ofTMS-238.� Both concept alloys, TMS-238MoRu andTMS-238MoIr, developed exceptionally fine c/c¢interfacial dislocation networks at 1100 �C/137 MPa, indicating a higher magnitude of negativelattice misfit. This was achieved through the additionof 1.5 at. pct Mo, which preferentially partitions intoc matrix. Suppression of TCP precipitates by Iraddition and improving lattice misfit strengtheningby Mo addition are the reasons for TMS-238MoIrexhibiting the best creep life at 1100 �C/137 MPa.� The underperformance of TMS-238MoRu com-pared to the baseline TMS-238 was more pro-nounced under lower temperature creep conditionsbetween 800 �C and 1000 �C. While TMS-238MoIrhas comparable creep life to TMS-238 at 1000 �C/245 MPa, this material also showed inferior creepperformance at 900 �C/392 MPa and 800 �C/735 MPa.� A high density of stacking faults was observed in thec matrix of TMS-238MoRu and TMS-238MoIr,even in the absence of external load. These faultslikely formed during the heat treatment due to thecombining effect of decreased stacking fault energyby Mo addition, and large coherency stress byexceptionally large lattice misfit. Such defects mayhave facilitated TCP nucleation at 900–1000 �C asthey lie on the same {111} crystallographic plane asTCP phases, and serve as segregation sites forTCP-forming solutes such as Re and W.� Post-creep microstructural analyses and Time-Tem-perature-Transformation (TTT) diagrams revealedthat Ir has a stronger TCP suppression effect thanRu. On the other hand, the detrimental effect ofexcessive Ru addition was evident, asTMS-238MoRu showed severe TCP precipitationduring creep tests at temperatures below 1000 �C,resulting in a creep rupture life shorter than thebaseline TMS-238. The TCP precipitates acted asdeformation sources during creep loading, andinitiated a chain of recrystallization from creepvoids, which subsequently served as crack propaga-tion paths.ACKNOWLEDGMENTSThe authors are grateful to Mariko Iguchi, KayokoNakakita, and Megumi Noro (Research Center forStructural Materials, NIMS) for technical assistancewith the sample preparation and microstructure obser-vations. A part of this work was supported by theNIMS Microstructural Characterization Platform(NMCP) as a program of the ‘‘Nanotechnology Plat-form’’ of the Ministry of Education, Culture, Sports,Science and Technology (MEXT), Japan. We aregrateful to N. Isaka, NMCP, NIMS, and K. Ogawa,Vibration Control Materials Group, NIMS, for techni-cal support in the TEM observation. This research wasfinancially supported by Japan Science and Technol-ogy (JST), under the Advanced Low Carbon Technol-ogy Research and Development Program (ALCA)project ‘‘Development of direct and complete recyclingmethod for superalloy turbine aerofoils.’’ Grant Num-ber JPNJAL1302.OPEN ACCESSThis article is licensed under a Creative CommonsAttribution 4.0 International License, which permitsuse, sharing, adaptation, distribution and reproductionin any medium or format, as long as you give appro-priate credit to the original author(s) and the source,provide a link to the Creative Commons licence, andindicate if changes were made. The images or otherthird party material in this article are included in thearticle’s Creative Commons licence, unless indicatedotherwise in a credit line to the material. 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Toshiharu Kobayashi, Yutaka Koizumi, Hideyuki Murakami,Yoshikazu Ro, Yoko Yamabe, Shizuo Nakazawa, Hiroshi Har-ada, Toshihiro Yamagata, Iridium-containing nickel-basesuperalloy, US6494971B1, 2002. https://patents.google.com/patent/US6494971B1/en.Publisher’s Note Springer Nature remains neutral with regard tojurisdictional claims in published maps and institutional affiliations.1812—VOLUME 57A, MAY 2026 METALLURGICAL AND MATERIALS TRANSACTIONS Ahttps://doi.org/10.1007/s11661-022-06713-2https://doi.org/10.1007/s11661-022-06713-2https://doi.org/10.1016/j.mtnano.2021.100152https://doi.org/10.1016/j.mtnano.2021.100152https://doi.org/10.1016/j.actamat.2025.121448https://doi.org/10.1016/j.actamat.2025.121448https://doi.org/10.1016/j.jallcom.2016.02.122https://doi.org/10.1016/j.jallcom.2016.02.122https://doi.org/10.1016/S1359-6454(03)00105-8https://doi.org/10.1007/s11661-018-4669-1https://doi.org/10.1007/s11661-018-4669-1https://doi.org/10.1016/j.actamat.2020.02.012https://doi.org/10.1016/j.actamat.2020.02.012https://doi.org/10.1007/s11661-024-07626-yhttps://doi.org/10.1016/j.mtnano.2025.100694https://doi.org/10.2320/jinstmet.69.1099https://doi.org/10.2320/jinstmet.69.1099https://furuyametal.jp/english/fmbi/chart/?language=enhttps://furuyametal.jp/english/fmbi/chart/?language=enhttps://doi.org/10.7449/2008/Superalloys_2008_499_508https://doi.org/10.7449/2008/Superalloys_2008_499_508https://patents.google.com/patent/US6494971B1/enhttps://patents.google.com/patent/US6494971B1/en Challenging the Limit of Strengthening by gamma / gamma vprime Lattice Misfit on the High Temperature Creep Properties of Ni-Base Single Crystal Superalloy Abstract Introduction Experimental Details Alloy Design Strategy and Single Crystal Sample Preparation High-Temperature Tests Microstructure Observations Results Microstructure Before the High-Temperature Tests Creep Properties of TMS-238 and Concept Variants Microstructural Analyses After Creep Tests at 1100 degC/137 MPa Microstructural Analyses After Creep Tests at 900 degC/392 MPa  gamma / gamma vprime Phase Stability and TCP Phase Precipitation Discussion Effect of Alloy Composition Modifications on the gamma / gamma vprime Lattice Misfit and the Creep Properties at 1100 degC/137 MPa Influence of Ru and Ir on the gamma / gamma vprime Phase Stability, TCP Formation, and Creep Properties of Ni-Base SC Superalloy Ni-Base SC Superalloy Design Including Ir Conclusion Open Access References