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[Fumiyoshi Yoshinaka](https://orcid.org/0000-0003-0534-7815), [Nobuo Nagashima](https://orcid.org/0000-0003-3588-980X), [Takahiro Sawaguchi](https://orcid.org/0000-0002-9405-002X)

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[Effect of Strain Rate on the Extremely Low-Cycle Fatigue of Fe-15Mn-10Cr-8Ni-4Si Bidirectional-TRIP Steel](https://mdr.nims.go.jp/datasets/b58f99a4-d023-45a5-b769-fc864da2204a)

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Effect of Strain Rate on the Extremely Low-Cycle Fatigue of Fe-15Mn-10Cr-8Ni-4Si Bidirectional-TRIP SteelEffect of Strain Rate on the Extremely Low-Cycle Fatigue of Fe-15Mn-10Cr-8Ni-4SiBidirectional-TRIP Steel+1Fumiyoshi Yoshinaka+2, Nobuo Nagashima and Takahiro SawaguchiNational Institute for Materials Science, Tsukuba 305-0047, JapanExtremely low cycle fatigue tests, up to a total axial strain amplitude of 10%, were conducted on Fe-15Mn-10Cr-8Ni-4Si bidirectionaltransformation-induced plasticity (B-TRIP) steel. The fatigue life was approximately five times longer than that of SUS316 when the total strainamplitude was 4% or higher. The improved fatigue life of Fe-15Mn-10Cr-8Ni-4Si was attributed to reversible bidirectional γ § εtransformation during fatigue deformation, which might mitigate fatigue damage. In contrast, the fatigue life tended to decrease with increasingstrain rate when the strain rate was varied from 0.1 to 2.5%/s with a total strain amplitude of 10%. Fractography revealed that the fracturesurface varied significantly with strain rate. At low strain rates, crystallographic fracture surfaces characterized by facets and secondary crackswere observed, whereas these features were not observed at high strain rates. Electron backscatter diffraction measurements of the postmortemmicrostructure showed that frequent ε-martensite formation occurred at low strain rates, whereas martensitic transformation was suppressed athigh strain rates. The change in the specimen surface temperature was evaluated in terms of the Gibbs free energy difference between γ-austeniteand ε-martensite (i.e., ¦Gγ¼ε), and the effect of strain rate on the extremely low cycle fatigue was discussed from the viewpoint of thedeformation mechanism. At a low strain rate, the condition for B-TRIP to work effectively, that is, ¦G γ¼ε is negative but close to zero, wasmaintained over the entire life span. At a high strain rate, the deformation mechanism changed to one in which γ-austenite was dominant becauseof the increase in ¦Gγ¼ε caused by self-heating; the fatigue damage mitigation mechanism provided by B-TRIP was less likely to occur at highstrain rates, which reduced the fatigue life. [doi:10.2320/matertrans.MT-Z2024002](Received February 15, 2024; Accepted April 16, 2024; Published April 26, 2024)Keywords: fatigue, fracture surface, extremely low cycle fatigue, microstructure, martensitic transformation, seismic damper1. IntroductionA vibration control structure is one in which seismicdampers act as energy-absorbing components to reduce theshaking of the building and minimize damage to the mainstructural members such as columns and beams. There areseveral types of seismic dampers, including steel, oil, andrubber dampers. Steel dampers absorb seismic energy inputto a building by converting it to heat through elastoplasticdeformation. Steel dampers offer advantages over other typesof dampers because of their low cost and ability to addstiffness to buildings. However, since their operationprinciple involves cyclic elastoplastic deformation, theyexhibit low cycle fatigue. Therefore, their design should bebased on the limited life design concept. In other words, theproduct life of a steel damper depends on the fatiguedurability against low cycle fatigue (or plastic fatigue) of thesteel used as the core material.Japan regularly experiences major earthquakes, such as theGreat East Japan Earthquake in 2011, the KumamotoEarthquake in 2016, the Hokkaido Eastern Iburi Earthquakein 2018, and the Noto Earthquake in 2024. Additionally, theNankai Trough Earthquake is predicted to occur with a highprobability. Therefore, earthquake preparedness is extremelyimportant. Recently, the problem of long-term long-periodground motion has attracted much attention. Structures thathave a long natural period, such as high-rise buildings, mayresonate with the long-period component of ground motion,causing long-lasting strong shaking. Conventional steeldampers cannot cope with long-term long-period groundmotion because the plastic fatigue life of conventional steel istoo short to endure the severe cyclic deformation.In general, the type of steel is unlikely to significantlyimpact plastic fatigue life [1]. However, our previous workdeveloped Fe-15Mn-10Cr-8Ni-4Si alloy, which exhibits along plastic fatigue life because of its deformationreversibility [2–5]. This Fe-15Mn-10Cr-8Ni-4Si alloy hasbeen used as a seismic damper in some structures in Japan.The deformation mechanism of this alloy is characterizedby reversible bidirectional transformation between the face-centered-cubic structure, γ-austenite, which is an initialmicrostructure, and the hexagonal-closed-packed structure,ε-martensite; this transformation is called γ § ε trans-formation [5]. Extended plastic fatigue life may beattributed to fatigue damage mitigation resulting from highdeformation reversibility.In practical applications of Fe-15Mn-10Cr-8Ni-4Si alloyas a seismic damper, the expected deformation, specifically,the total axial strain amplitude, is up to 1% during shakingcaused by L2 ground motion under extremely rare earth-quakes. When a joint part is subjected to fatigue deformation,it is necessary to consider the occurrence of significant strainconcentration, and when a damper shape such as a U shape[6] or a lens shape [7] is used, it is necessary to assume thatlarge deformation will occur repeatedly. In addition, it isrequired to consider L3 ground motion under extremeearthquakes that exceed the conventional assumption of L2ground motion. Therefore, it is crucial to understand theextremely low cycle fatigue property of Fe-15Mn-10Cr-8Ni-4Si alloy to guarantee its integrity when used as a seismicdamper.Our previous work conducted extremely low cycle fatiguetests on the alloy involving a maximum total strain amplitudeof 10%, converted to the axial direction from the radialdirection [8]. The present work conducted fatigue tests at a+1This Paper was Originally Published in Japanese in J. Japan Inst. Met.Mater. 72 (2023) 858–865. The captions of Table 1 and Figs. 1, 6, and 7are slightly modified.+2Corresponding author, E-mail: YOSHINAKA.Fumiyoshi@nims.go.jpMaterials Transactions, Vol. 65, No. 7 (2024) pp. 773 to 779©2024 The Society of Materials Science, Japanhttps://doi.org/10.2320/matertrans.MT-Z2024002total axial strain amplitude of 10% with different strain rates.The effect of strain rate on the extremely low cycle fatigueof Fe-15Mn-10Cr-8Ni-4Si alloy was determined based onfracture surface observation and microstructural analysis.2. Experimental ProcedureAn Fe-15Mn-10Cr-8Ni-4Si (mass%) alloy was used in thisstudy. The microstructure after heat treatment at 1273K for1 h followed by water quenching was γ-austenite with anaverage grain size of 95 µm.A radial strain-controlled fatigue test was conducted usinga 50 kN servo-hydraulic fatigue testing machine. Triangularwaves were applied to the specimen at a strain ratio of ¹1,total radial strain amplitudes (¾td, a) of 2, 3, 4, and 5%, and anaxial strain rate of 0.5%/s (hereafter, these tests are referredto as the single strain rate tests). In addition, tests at ¾td, a =5% were performed with axial strain rates of 0.1, 0.5, and2.5%/s to study the effect of strain rate (hereafter, these testsare referred to as the multiple strain rate tests). An hourglass-shaped specimen with a minimum diameter of 6mm wassubjected to the single strain rate tests, and an hourglass-shaped specimen with a minimum diameter of 8mm wassubjected to the multiple strain rate tests. Radial strain wasmonitored to control the test using an extensometer attachedto the minimum diameter part along the radial direction. Thetemperature at the specimen surface was measured duringthe fatigue test using a Type T thermocouple wire. Themeasurement point was approximately centered in thespecimen. The radial strain ¾td was converted to axial strain¾tl by the following equation:¾tl, a ¼ ð·a=EÞð1� 2¯Þ þ 2¾pd, a ð1Þwhere ¾tl, a, ·a, E, and ¯ are the total axial strain amplitude,stress amplitude, Young’s modulus (189GPa), and Poisson’sratio (0.27), respectively. The total strain amplitude can bewritten as the summation of the elastic and plasticcomponents, and the first term on the right-hand side ofeq. (1) shows the elastic component. The second termrepresents the plastic component, where ¾pd, a was measuredas half of the distance between the x-axis (strain axis)intercepts on the hysteresis loop, i.e., half of the plastic radialstrain range. Hereafter, unless otherwise noted, strainamplitude and strain rate refer to axial values.Fractography was performed using a field-emissionscanning electron microscope (FE-SEM) (Thermo Fisher,Scios2), and Thermo Scientific MAPS software was used toautomatically capture the overall fracture surface.The fracture surface was cut parallel to the direction of theload axis using a precision cutting machine to expose thelongitudinal section for analysis of the microstructure beneaththe fracture surface. The cutting positions were approx-imately through the fracture origin. The sample was wetpolished with #180–#1200 waterproof abrasive paper,followed by intermediate polishing with 9, 3, and 1 µmdiamond abrasives, and then finished with colloidal silicafor a mirror finish. Observations were performed using anFE-SEM (JEOL, JSM-7900F), and electron backscatterdiffraction (EBSD) measurements were conducted using adetector (DVC5, TSL) equipped with the FE-SEM. TheEBSD data were analyzed using OIM Analysis v8 software(EDAX).3. ResultsFigure 1 shows the ¾-N curve, representing the relationshipbetween the total strain amplitude ¾tl, a and the fatigue life Nfobtained by the single strain rate tests (0.5%/s) [8]. Theresults of the axial strain-controlled fatigue tests at total strainamplitudes of 0.25%–2% obtained with a strain rate of 0.4%/s in the previous work ( ) [3] and the results of the multiplestrain rate tests in the present work ( ) are plotted in thefigure along with the results of the single strain rate tests ( ).Additionally, the results of SUS316 obtained by Kamaya ( )[9] are presented in the figure for comparison. The Fe-15Mn-10Cr-8Ni-4Si alloy exhibited an extended fatigue lifecompared to SUS316. In the regime of extremely low cyclefatigue, the fatigue life of Fe-15Mn-10Cr-8Ni-4Si alloy wasapproximately five times longer than that of SUS316, and itwas found to be 28 cycles at ¾tl, a = 10%.Table 1 lists the fatigue lives at ¾tl, a = 10% under strainrates of 0.1, 0.5, and 2.5%/s obtained by the multiple strainrate tests. As shown in the table, fatigue life decreased withincreasing strain rate. The fatigue life of SUS316 at ¾tl, a =10%, estimated by extrapolating the results shown in Fig. 1,was Nf = 9 cycles. Thus, although the fatigue life of Fe-15Mn-10Cr-8Ni-4Si alloy decreased with increasing strainrate, it remained superior (approximately two times longer)compared to that of conventional steel.Figure 2 shows the relationship between N and themaximum stress ·max from the multiple strain rate tests.100 101 102 103 104 105 106Fatigue life Nf (cycles)0.1110Total strain amplitude ε t, a(%)  : Fe-15Mn-10Cr-8Ni-4Si (this work): Fe-15Mn-10Cr-8Ni-4Si (previous work): SUS316: 2.5%/s: 0.5%/s: 0.1%/sFig. 1 Relationship between the total strain amplitude and fatigue life.Table 1 Results of fatigue tests at a total axial strain amplitude of 10% anddifferent strain rates.F. Yoshinaka, N. Nagashima and T. Sawaguchi774Remarkable hardening occurred during the initial part ofthe fatigue tests at all strain rates. In the test at 0.1%/s, thehardening became saturated with additional cyclic loadings,whereas in the tests at 0.5 and 2.5%/s, the hardeningcontinued until fatigue fractures occurred. Significant hard-ening was observed, particularly at 2.5%/s. The maximumstress at half-life Nf/2 increased with the strain rate, as shownin Table 1.Figure 3 shows the specimen temperature during thefatigue tests. In the test at 0.1%/s, the temperature wasalmost stable at room temperature, whereas in the tests at 0.5and 2.5%/s, the temperature increased with cyclic loadings.Significant self-heating was observed, particularly at 2.5%/s.In the following section, the mechanism of extremely lowcycle fatigue of Fe-15Mn-10Cr-8Ni-4Si alloy is investigatedby analyzing the fatigue fracture and deformation micro-structure obtained from the multiple strain rate tests.Figures 4(a)–(c) show the overall fracture surfaces. Thefracture origin was on the specimen surface regardless ofthe strain rate. Figures 4(d) and (e) show the magnified viewof the region formed by fatigue crack growth in the tests at101 102100Number of cycles N (cycles)900800700600500Maximum stress σmax[MPa]2.5%/s 0.5%/s0.1%/sFig. 2 Relationship between the number of cycles N and the maximumstress ·max at total axial strain amplitude of 10%.605040302070Specimen temperature T[℃]101 102Number of cycles N (cycles)1002.5%/s0.5%/s0.1%/sFig. 3 Relationship between the number of cycles N and the specimentemperature T during the fatigue test at a total axial strain amplitude of10%.Fig. 4 Fracture surface. (a)–(c) Overall fracture surface obtained with strain rates of 0.1%/s, 0.5%/s, and 2.5%/s. (d)–(e) Fatigue crackgrowth region in (a) and (c). (f )–(h) Magnified view of the fracture surface at the fracture origin, the fatigue crack growth region, andthe fast fracture region in (a).Effect of Strain Rate on the Extremely Low-Cycle Fatigue of Fe-15Mn-10Cr-8Ni-4Si Bidirectional-TRIP Steel 7750.1%/s and 2.5%/s. The fracture surface obtained at 0.1%/shad an angular crystallographic morphology, frequentlydisplaying facets and secondary cracks. In contrast, thefracture surface obtained at 2.5%/s had a roundedmorphology without facets. Figures 4(f )–(h) present exam-ples of the crystallographic morphologies observed on thefracture surface obtained at 0.1%/s. Facets and secondarycracks were observed in the origin site and the fatigue crackgrowth region. Secondary cracks were often found at theedges of facets. In the final fracture region after fatigue crackgrowth, dimples coexisting with facets were observed. In thecase of 2.5%/s strain rate, only dimples were observed in thefinal fracture region. The fatigue fracture surface significantlychanged with varying strain rate, as shown in Fig. 4, and acrystallographic fracture surface characterized by facets andsecondary cracks was formed when the strain rate was low.Figures 5(a) and (b) show the SEM images of thelongitudinal section obtained at strain rates of 0.1%/s and2.5%/s, respectively. The macrographic crack growthdirection was oriented from right to left. The blue openarrows, green blank closed arrows, and red solid closedarrows point at secondary cracks from the main crack, surfacecracks, and internal cracks, respectively. Although thefracture origin was on the specimen surface, as shown inFig. 4, internal cracks were also present on the longitudinalsections. The number of cracks on the longitudinal sectionvaried significantly depending on the strain rate, with manycracks observed when the strain rate was 0.1%/s.EBSD analysis was conducted on longitudinal sections.The measurement areas located beneath the fatigue crackgrowth regime are represented by the dashed rectangles inFigs. 5(a) and (b). The size of the measurement area was800 µm © 800 µm, and the measurement step was 1 µm. TheEBSD phase and grain boundary maps are shown inFigs. 6(a) and (b) for 0.1%/s and Figs. 6(c) and (d) for2.5%/s. The figures exclude measurement points with aconfidential index value less than 0.1. The phase maps aresuperimposed on the image quality maps. A significantamount of ε-martensite was formed when the strain ratewas 0.1%/s, as shown in Fig. 6(a), and the retained γ-austenite locally existed. In addition, limited amounts ofαA-martensite were detected. As shown in Fig. 6(c), thedeformation microstructure at 2.5%/s consisted only ofγ-austenite. Deformation twinning was only found at 2.5%/s,as shown in the grain boundary maps. Deformation micro-structure varied with strain rate. At low strain rates, the ε-martensitic transformation occurred frequently, whereas athigh strain rates, the transformation was suppressed, and thedeformation mechanism changed to one in which γ-austenitewas dominant.4. DiscussionFe-15Mn-10Cr-8Ni-4Si alloy had an extended fatigue lifeeven at strain levels that cause conventional steels to exhibitextremely low cycle fatigue, as shown in Fig. 1. The alloyunderwent bidirectional γ § ε transformation as a plasticdeformation mechanism, referred to as B-TRIP [10–13].High deformation reversibility can improve fatigue life [14–16]. Fe-15Mn-10Cr-8Ni-4Si alloy possesses high deforma-tion reversibility in the form of bidirectional transformation,resulting in extended fatigue life. Although its deformationFig. 5 Longitudinal section at the fatigue crack propagation regime: strain rate of (a) 0.1%/s and (b) 2.5%/s.Fig. 6 Electron backscattering diffraction maps of the postmortem microstructure obtained from the region shown in Fig. 4: (a) phase and(b) grain boundary maps for strain rate of 0.1%/s; (c) phase and (d) grain boundary maps for a strain rate of 2.5%/s.F. Yoshinaka, N. Nagashima and T. Sawaguchi776reversibility decreases with increasing strain, Fe-15Mn-10Cr-8Ni-4Si alloy still exhibits superior fatigue life comparedto conventional steel at a high strain level [15]. Therefore,the fatigue improvement mechanism via B-TRIP may beeffective in the extremely low cycle fatigue regime.The fatigue life and cyclic hardening behavior of Fe-15Mn-10Cr-8Ni-4Si alloy varied with strain rate, as shownin Table 1, respectively. Furthermore, the fatigue fracturesurface and deformation microstructure were dependent onthe strain rate. As shown in Fig. 4, the fatigue fracturesurfaces obtained at a low strain rate (resulting in longerfatigue life) often displayed facets and secondary cracks,whereas these features were not observed in the fracturesurface obtained at a high strain rate (resulting in shorterfatigue life). In terms of the fatigue deformation micro-structure, large amounts of ε-martensite were detected at alow strain rate, whereas ε-martensitic transformation wassignificantly suppressed at a high strain rate. Previous workon the fracture surface of Fe-33Mn-6Si alloy reported that ithad a crystallographic appearance consisting mainly of facetswith frequent secondary cracks [17]. This may be attributedto the high probability of ε-martensitic transformation dueto its low stacking fault energy (SFE) !SFE. Therefore, thecrystallographic fracture surface observed in Fe-15Mn-10Cr-8Ni-4Si alloy at a low strain rate may also be attributed toε-martensite. The absence of facets on the fracture surfaceat a high strain rate can be explained by the significantsuppression of ε-martensite.Austenitic steels, including Fe-15Mn-10Cr-8Ni-4Si alloy,exhibit different plastic deformation mechanisms dependingon the SFE [18].� SFE ¼ 2µ�G�!" þ 2·�=" ð2Þwhere the molar surface fraction on the γ plane is µ, the Gibbsfree energy difference between γ-austenite and ε-martensiteis ¦Gγ¼ε, and the surface entropy interfacial surface energybetween the γ- and ε-phases is ·γ/ε. A detailed descriptionof the !SFE and ¦Gγ¼ε is presented in the previous paper[19]. The plastic deformation mechanism changes fromstress-assisted ε-martensitic transformation to strain-inducedε-martensitic transformation, deformation γ-twinning, and γ-slip with increasing SFE. For instance, Fe-33Mn-6Si alloycan readily undergo ε-martensitic transformation due to itslow SFE.Previous work on the low cycle fatigue behavior of Fe-Mn-Al-Si alloys with systematically varied SFE demon-strated that the cyclic hardening behavior varied accordingly[20]. Significant cyclic hardening occurred in all alloys.However, the lower SFE alloys exhibited saturation inhardening with the progress of the fatigue cycle, whereasthe higher SFE alloys showed noticeable secondary harden-ing. A previous study proposed that a higher SFE, whichmakes cross slip more likely to occur, triggers an interactionwith the lattice dislocation, resulting in secondary hardening.As shown in Fig. 2, Fe-15Mn-10Cr-8Ni-4Si alloy exhibitedsaturation in the hardening when it had lower SFE, whereasit showed secondary hardening when it had higher SFE. Thehigh strain rate test demonstrated significant self-heating,leading to an increase in SFE. In general, a high strain ratecauses an increase in the deforming stress due to enhancedviscous resistance to dislocation motion. However, theincrease in SFE due to self-heating may play a role in thestrain rate dependence of the cyclic hardening behavior inthe Fe-15Mn-10Cr-8Ni-4Si alloy.In our previous work, the occurrence conditions for B-TRIP were expressed as ¦Gγ¼ε ¯ 0 (¦Gγ¼ε is negativeand close to zero) [21]. Additionally, the relationship betweenthe plastic deformation mechanism and the fatigue life wascharacterized by ¦Gγ¼ε [19]. Therefore, in this study,¦Gγ¼ε was used instead of !SFE to discuss the effect ofstrain rate. The increase in ¦Gγ¼ε, which has a positivedependence on temperature, implies an increase in thestability of γ-austenite with respect to ε-martensite. There-fore, the significant strain rate dependence of the fatiguelife of Fe-15Mn-10Cr-8Ni-4Si, shown in Table 1, may beattributed to the increase in ¦Gγ¼ε due to self-heating at ahigher strain rate. Thus, the specimen temperature shown inFig. 3 was converted to ¦Gγ¼ε, and the change in ¦Gγ¼εwith the number of cycles is shown in Fig. 7. As shown inFigs. 3 and 7, no significant self-heating occurred at a strainrate of 0.1%/s, and the condition of ¦Gγ¼ε ¯ 0 was satisfiedthroughout the fatigue life, resulting in enhanced fatiguedurability due to B-TRIP. As shown in Fig. 6(a), thepostmortem microstructure was mainly occupied by ε-martensite at the strain rate of 0.1%/s. As previouslymentioned, Fe-33Mn-6Si alloy can readily undergo ε-martensitic transformation, resulting in a large amount ofε-martensite after fatigue fracture. An extremely low cyclefatigue test has not yet been conducted on Fe-33Mn-6Sialloy, but its fatigue life at a total strain amplitude of less than2% was not long [17]. Therefore, the ε-martensitic trans-formation is not likely to contribute to the improvementof fatigue life. The Fe-15Mn-10Cr-8Ni-4Si alloy exhibitedenhanced fatigue life due to B-TRIP; however, not all of theε-martensite formed by the forward transformation (γ ¼ εtransformation) underwent the reverse transformation (ε¼ γtransformation), and the gradual accumulation of ε-martensiteup to the fatigue fracture may have developed thedeformation microstructure shown in Fig. 6(a).As shown in Fig. 5, many cracks were initiated inside thematerial at a low strain rate. It has been reported that thefracture mode caused by internal cracks due to the staticfracture mechanism occurs in the extremely low cycle fatiguein addition to the fracture mode caused by surface cracks101 1021002.5%/s0.5%/s0.1%/s-80-60-40-2004020Gibbs free energy difference between γand εΔ Gγ →ε[mJ/m2 ]Number of cycles N (cycles)Fig. 7 Relationship between the number of cycles N and the Gibbs freeenergy difference between γ-austenite and ε-martensite, that is, ¦Gγ¼ε,during fatigue test at a total axial strain amplitude of 10%.Effect of Strain Rate on the Extremely Low-Cycle Fatigue of Fe-15Mn-10Cr-8Ni-4Si Bidirectional-TRIP Steel 777according to the general fatigue crack growth law [22].Koyama et al. reported that the formation of ε-martensitetriggers the brittle-like crack initiation, resulting in prematurefracture in the tensile fracture of high-Mn steel [23].Therefore, many internal cracks found at the low strain ratedespite the extended fatigue life can be attributed to thedeformation-induced ε-martensite, which can be the initiationsite of the crack due to the static fracture mechanism.However, the Fe-15Mn-10Cr-8Ni-4Si alloy showed surfacefracture regardless of the strain rate, and there was noevidence of coalescence with internal cracks initiatedindependently of the main (surface) crack on the fracturesurface, as shown in Fig. 4. Therefore, the fatigue life of thisalloy can be evaluated using the fatigue crack growth law forgeneral low cycle fatigue, at least in the range of strain ratesinvestigated in the present work.At a strain rate of 2.5%/s, the ¦Gγ¼ε increased andeventually became positive as the test progressed, resulting inthe deformation of γ-austenite without the martensitictransformation, as shown in Fig. 6(c). Significant self-heatingdue to high strain rate, which increased ¦Gγ¼ε, inhibitedthe bidirectional transformation, resulting in a decrease infatigue life. However, the Md point—the upper limit of thetemperature at which martensite can be induced by plasticdeformation—of Fe-15Mn-10Cr-8Ni-4Si alloy is 100°C [19],and the maximum temperature measured at a strain rate of2.5%/s (shown in Fig. 3) was lower than the Md point.The measurement point of the specimen temperature wasapproximately located at the minimum diameter, but it wasimpossible to place the thermocouple at the exact point atwhich maximum strain occurred because the extensometerneeded to be placed there. Self-heating is expected to be morepronounced in the locations at which large deformationsoccur, and the temperature at the minimum diameter mayhave been higher than that shown in Fig. 3. In addition, thefatigue crack is expected to cause further self-heating at itstip [24] thereby increasing SFE. Figure 5 was taken fromthe lower region of the fracture surface, and as shown inFig. 6(c), little martensitic transformation occurred at highstrain rates, which can be explained by the local heatgeneration near the crack. A detailed analysis of the localheat generation at the crack propagation zone and the fatiguedeformation microstructure around the crack is a subject forfuture study.As described in Section 3, the fatigue life of Fe-15Mn-10Cr-8Ni-4Si alloy decreased with increasing strain rate andwas 25 cycles, which was more than double that of SUS316,even at the highest strain rate of 2.5%/s. The decrease infatigue life was due to self-heating, as discussed above.However, the increase in specimen temperature due to self-heating gradually occurred over the entire fatigue life at ahigh strain rate (Fig. 3), as did the associated increase in¦Gγ¼ε (Fig. 7). Therefore, B-TRIP was effective in theearly stage of fatigue life, leading to the longer fatigue lifecompared to general steel at a high strain rate. In addition, itis unrealistic to assume that high-speed large deformationswill continue until fatigue fracture during a real-worldearthquake. Instead, intermittent deformations are more likelyto occur. In this case, even if self-heating occurs because ofhigh-speed large deformations, if sufficient heat removal isachieved, B-TRIP in the subsequent fatigue deformation andthe fatigue damage accumulation mechanism might beeffective. In other words, the fatigue test under constant highstrain rate conducted in the present work is significantbecause it represents the worst case of fatigue fracture ofFe-15Mn-10Cr-8Ni-4Si alloy. In conclusion, Fe-15Mn-10Cr-8Ni-4Si alloy exhibits excellent fatigue life under large cyclicdeformation when the heat generated during deformation isproperly controlled. This alloy helps to ensure designmargins against extreme earthquakes and has strong potentialfor application in various types of seismic dampers.5. ConclusionFe-15Mn-10Cr-8Ni-4Si alloy, developed for steel damperexhibits bidirectional γ § ε transformation (B-TRIP) as aplastic deformation mechanism that mitigates the accumu-lation of fatigue damage, resulting in approximately fivetimes longer fatigue life compared to SUS316 at a total axialstrain amplitude over 4%. However, the fatigue life decreasedwith increasing strain rate between 0.1%/s and 2.5%/s at atotal axial strain amplitude of 10%. The fatigue fracturesurface also changed with strain rate; facets and secondarycracks were observed at a low strain rate (longer fatigue life),whereas these features were hardly observed at a high strainrate (shorter fatigue life). EBSD analysis of the deformationmicrostructure demonstrated that the γ ¼ ε-martensitic trans-formation occurred frequently at a low strain rate, whereasthe transformation was suppressed at a high strain rate. Thespecimen temperature measurement revealed that significantself-heating occurred only at a high strain rate of 2.5%/s,indicating that the increase in ¦Gγ¼ε due to self-heatinginhibited B-TRIP, resulting in a decrease in fatigue life athigh strain rates. The fact that the Fe-15Mn-10Cr-8Ni-4Sialloy maintained a relatively long fatigue life compared togeneral steel implies that the fatigue damage mitigationmechanism due to B-TRIP was effective at the early stage offatigue life when less self-heating had occurred.AcknowledgmentsThe authors acknowledge the support of Grants-in-Aidfor Early-Career Scientists (19K14853 and 23K13227) and aGrant-in-Aid for Scientific Research (C) (20K04170) fromthe Japan Society for the Promotion of Science, Japan.REFERENCES[1] K. 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