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Materials Science & Engineering A

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Effect of austenitic texture on tensile behavior of lean duplex stainless steelwith transformation induced plasticity (TRIP)

Jun-Yun Kanga,⁎, Hoyoung Kima, Kyung-Il Kimb, Chang-Hoon Leea, Heung Nam Hanb,Kyu-Hwan Ohb, Tae-Ho Leea

a Korea Institute of Materials Science, 797 Changwon-daero, Seongsan-gu, Changwon, Gyeongnam 51508, Republic of Koreab Department of Materials Science and Engineering and Research Institute of Advanced Materials, Seoul National University, 1 Gwanak-ro, Gwanak-gu,Seoul 08826, Republic of Korea

A R T I C L E I N F O

Keywords:SteelAusteniteMechanical characterizationMartensitic transformationTexture

A B S T R A C T

Mechanically induced martensitic transformation (MIMT) and consequent plastic flow behavior with respect toaustenitic texture were investigated in a lean duplex stainless steel. Different grain sizes and textures with fixedphase fractions were obtained via varying the thermomechanical processes. Nearly random distribution ofaustenitic orientation exhibited a distinguished flow curve from the others with a major D {4 4 11}⟨11 11 8⟩component due to more gradual enhancement of hardening by less martensitic transformation. In order tocompare the susceptibility to the transformation with respect to individual austenitic orientations and theexperimental textures, interaction energy between the imposed stress and transformation strain was calculatedby a classical transformation and a crystal plasticity model. The results indicated that a larger stress imposed onthe D component led to higher interaction energy and a steeper progress of MIMT observed in the texturedmaterials.

1. Introduction

Duplex stainless steels (DSS) have microstructures that consist ofsimilar fractions of face centered cubic (fcc) austenite and bodycentered cubic (bcc) ferrite [1]. By virtue of constructive propertycombinations between the two constituent phases, they usually showan excellent balance in mechanical properties as well as corrosionresistance. They are used in many applications in chemical, petroleum,and atomic energy industries [2–5], and have progressively substitutedfor some classical austenitic stainless steels since the 1990s [6].

Lean DSSs have been developed to reduce the initial material costin production and cost instability by reduction of expensive alloyelements [6–11]. As a major strategy to design lean compositions,inexpensive austenite stabilizers such as N and Mn have been added tolower the content of an expensive one, i.e., Ni [6–11]. In addition to theeffect of austenite stabilization, Mn increases the solubility of theinterstitial element N [12,13] which gives potent solid solutionstrengthening [13,14] and improves the resistance to pitting corrosion[13,15]. Therefore, in spite of the leaner compositions, some lean alloysreported superior mechanical properties and corrosion resistance [9–11].

Recently, a few lean DSSs with exceptionally good tensile properties

by transformation induced plasticity (TRIP) were introduced [16–21].In these alloys, the mechanically induced martensitic transformation(MIMT) in metastable austenite enhanced their strain hardeningcapacity, and the extended progress of MIMT to a large strain resultedin a very high tensile strength and ductility over 1 GPa and 60%respectively [18]. From the consideration on N content and itspartitioning between the phases, stacking fault energy (SFE) whichdetermined the deformation mechanisms of austenite was estimated tobe in the range for MIMT [19,20]. Detailed microscopic analyses wereperformed to interpret the relationship between the progress of MIMTand tensile behavior, which confirmed the beneficial effect of MIMT onthe excellent plasticity [20,21]. A number of studies followed to revealthe effect of various factors on this TRIP effect [21–24]. It is knownthat the strength and the ductility increased with an increasing contentof N because of the increasing volume of mechanically inducedmartensite [21]. It was also reported that enhanced MIMT below roomtemperature increased uniform elongation [22], while an increasedstrain rate [23] and annealing temperature [24] suppressed MIMT.However, in spite of these studies, less focus was put on the probableeffect of grain size, orientation or texture. In this article, the effect ofaustenitic texture on characteristic tensile behavior by TRIP is analyzedand discussed using a TRIP-aided lean DSS, which additionally

http://dx.doi.org/10.1016/j.msea.2016.11.001Received 13 June 2016; Received in revised form 31 October 2016; Accepted 1 November 2016

⁎ Corresponding author.E-mail address: firice@kims.re.kr (J.-Y. Kang).

Materials Science & Engineering A 681 (2017) 114–120

0921-5093/ © 2016 Elsevier B.V. All rights reserved.Available online 03 November 2016

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contributes to the development of lean DSSs with tailored properties.

2. Experimental

The composition of the alloy is presented in Table 1. An ingot of9 kg weight was cast using vacuum induction melting. It was reheatedto 1200 °C, held for 2 h, hot rolled between 1100 and 950 °C andquenched in water. The thickness reduction in the hot rolling was82.5% (from 40 mm to 7 mm). A part of the hot band was annealed at1050 °C for 30 min, quenched in water, and denoted as HR30. Theremaining parts were cold rolled by a thickness reduction of 82% (from7 to 1.25 mm), also annealed at 1050 °C for 5 or 120 min andquenched. The former was denoted as CR5 and the latter as CR120.The designations of the above specimens simply represent the processprior to the annealing, i.e., hot rolling (HR) or cold rolling (CR), andthe durations of the final annealing at 1050 °C, i.e., 5, 30 or 120 min.

Uniaxial tensile tests of the specimens were conducted at roomtemperature using a universal test machine (Instron 5882) with acrosshead speed of 2 mm/min, i.e., an initial strain rate of1.33×10−3 s−1. The preparation of the specimens and the test proce-dure followed the instructions in ASTM E8 [25]. The microstructures ofthe specimens were characterized using an electron backscatter dif-fraction (EBSD) system, Oxford Instruments NordlysNano detectorwith AZTEC software in a field emission scanning electron microscope(FESEM), JEOL JSM-7001F. EBSD mappings in the mid-thicknessregions of the specimens were conducted on electropolished surfacesthat were normal to the transverse direction (TD). This microtextureanalysis presented the morphological characteristics of grains, frac-tions and textures of the two constituent phases, ferrite (α) andaustenite (γ). The electropolishing was carried out using a commercialelectropolisher (Struers LectroPol5) with a solution of 90 vol% ethanoland 10 vol% perchloric acid at −20 °C. For statistical reliability in theanalyses, the total mapping area per specimen covered 0.8–3.2 mm2

depending on the grain size (at least 10,000 grains per phase).

3. Results

The microstructures of the specimens are presented in Fig. 1 via theoverlay images of phase, boundary and band contrast maps constructedfrom the EBSD mappings. The clustering of each phase and theconsequent banded structures were observed. The average grain sizesand phase fractions are listed in Table 2. All the specimens wereconstituted by nearly the same phase fractions due to the sameannealing temperature. It was clear that the minimum duration ofannealing, i.e., 5 min for CR5, was sufficient to attain the fully annealedand equilibriated microstructure because of the applied high tempera-

ture (1050 °C). Considerable grain growth occurred in the prolongedannealing of CR120 in which the average grain diameter increased tomore than three times from that of CR5. At a brief glance, HR30exhibited the largest grain size in Fig. 1 because of some coarse grains.However, its average grain sizes were comparable to those of CR120 inTable 2. The broader grain size distribution in HR30 could beattributed to the omission of cold rolling, while the ratio of itsmaximum grain size to that of CR120 was less than 3 (80 µm inHR30 vs. 30 µm in CR120) which was less than the average grain sizeratio of CR120 to CR5.

Fig. 2(a) presents the tensile flow curves, and the average tensileproperties are listed in Table 3. As expected from the grain sizedistributions in Fig. 1 and Table 2, CR5 had the highest stressevolution, followed by CR120. All the specimens exhibited very hightensile strength of over 800 MPa with large elongations exceeding 60%.Characteristic upward deflections of the flow curves were observed,which were more definite for CR5 and CR120 (indicated by the brokencircle). The less definite deflection of HR30 led to a distinct shape of thecurve, and the dissimilarity increased with strain (indicated by adouble-sided arrow). Fig. 2(b) corresponds to the strain hardeningrate obtained from Fig. 2(a). All the specimens experienced a transientincrease in the hardening rate after their respective local minimumvalues, σb (marked with arrows). Fig. 2(c) is the normalization ofFig. 2(b) with the respective σb. It shows more clearly that HR30exhibited later, slower and more gradual enhancement in hardening.This distinctive hardening behavior would explain the distinct shape ofthe flow curve as well as the increasing dissimilarity with strain.

Figs. 3 and 4 present the orientation distribution function (ODF) ofaustenite and ferrite, respectively, in the φ2=45° section of the Eulerspace (notation by Bunge [26]), in which most of major texturecomponents could be shown. For convenience, the major componentsfound in this study are listed in Table 4. CR5 and CR120 had almostidentical austenitic textures of moderate intensity as shown in Fig. 3.They were made up primarily of the D component with a very weakGoss component. As shown in Fig. 3(c), the peak ODF value for HR30was too small, thus, the austenitic orientation distribution could beregarded as random. In Fig. 4, RC (rotated cube) was the commonprimary component of ferritic textures for all the specimens. CR5 andCR120 show the typical types of textures observed in many cold rolledand annealed ferrtic steels [27,28]. While the ferritic texture of CR5more resembled cold rolling textures which were represented as astrong partial RD//⟨110⟩ fiber component, that of CR120 was closer toannealing textures with weakened RD//⟨110⟩ and enhanced ND//

Table 1Chemical composition of the lean duplex stainless steel (wt%).

Cr Mn Mo W C N Si

18.28 5.68 2.38 0.39 0.035 0.300 0.11

GB / PB

γγ

ND

RD

(b) (c)

100 μm

(a)

100 μm 100 μm

α

TB

Fig. 1. Microstructures according to the process conditions (overlay of phase, band contrast and boundary maps from EBSD): (a) CR5, (b) CR120, (c) HR30 (γ: austenite, α: ferrite, GB:grain boundary whose disorientation exceeds 3°, PB: phase boundary, TB: twin boundary, ND: normal direction, RD: rolling direction).

Table 2Average phase fractions and grain sizes in equivalent circle diameter (ECD).

Phase CR5 CR120 HR30

α Fraction (%) 41.1 42.4 42.5Grain size (μm) 2.72 9.82 9.10

γ Fraction (%) 58.9 57.6 57.5Grain size (μm) 3.31 10.27 13.54

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⟨111⟩. In spite of the similar features with cold rolling textures offerrite, it should be noted again that CR5 was fully annealed as alreadyshown in Fig. 1(a). HR30 had somewhat distinct ferritic textures withmore developed ND//⟨001⟩ fiber, while the maxima were still locatedat RC.

4. Discussion

4.1. Plastic flow behavior with respect to grain size and texture

As shown in Fig. 2(a), CR5 and CR120 with identical austenitictexture had almost parallel flow curves after about 10% of tensilestrain. Thus, it is clear that the differences in grain size and ferritictexture should have little influence on the shape of the flow curve. Thehardening rate curves in Fig. 2(b) were very similar, and the normal-ized ones in Fig. 2(c) even overlapped until their concurrent onset ofthe transient enhancement in hardening. The small difference in stresslevel and hardening behavior would more likely be due to thesubstantial grain size difference (more than three times on average)rather than to the relatively less remarkable difference in ferritictexture.

With the above consideration on the limited role of grain size and

ferritic texture, we thought that a comparison between CR120 andHR30 would be interesting as they have apparently different austenitictextures. Indeed, HR30 exhibited a distinctive shape of the flow curvein Fig. 2(a). The lower stress level of HR30 in Fig. 2(a) may be acceptedfrom the existence of some coarse grains in Fig. 1(c). However, itshould be noted that the difference in grain size was less substantialthan that of CR5 and CR120, as was stated in the previous section. Thedifferent ferritic texture could partly contribute to this lowered flowstress, while its contribution to the distinct shape of the flow curve wasstill questionable from the above comparison between CR5 and CR120.Therefore, it is reasonable to point out the last factor, i.e., the differentaustenitic texture as the major source for the distinct flow curve ofHR30.

The deformed microstructures in Fig. 5 provide a clearer insightinto the reason for the different plastic flow behaviors of CR120 andHR30. In both specimens, a large fraction of original austenites wastransformed to martensite by the tensile test. Thus, the transientenhancement of hardening and the resulting upward deflections of flowcurves in Fig. 2 can be attributed to this characteristic deformationbehavior of austenite, i.e., the progress of MIMT. It is generally knownthat similar transient enhancement in hardening is a typical sign forthe operation of MIMT which leads to enhanced plasticity by TRIP[18–20,29]. In Fig. 5(a) larger fraction of austenite was retained inHR30 from the nearly identical initial fractions (i.e., approximately58% in Table 2). The average remaining austenitic fractions in theuniformly strained volume were 15.0 and 26.2% respectively in CR120and HR30, which is equivalent to the less extensive progress of MIMTin HR30 in spite of a larger strain to flow localization and fracture. Asalready mentioned, the key characteristics of HR30 were the delayedand less intense transient hardening in Fig. 2(b) and (c) whichcontinuously enlarge the dissimilarity from CR120 in Fig. 2(a), morenotably in the large strain regime ( > 20%). This must have originated

Fig. 2. Uniaxial tensile behaviors according to the heat treatments: (a) engineering stress-strain curves (note that the break on the vertical axis for better differentiation of the curves),(b) hardening rate (dσ/dε) curves, (c) normalization of (b) by the local minima (σb).

Table 3Tensile properties (YS: yield strength evaluated with 0.2% proof stress, TS: ultimatetensile strength, UE: uniform elongation, TE: total elongation).

CR5 CR120 HR30

YS (MPa) 572.9 516.2 491.7TS (MPa) 952.5 929.8 850.0UE (%) 59.3 61.7 65.7TE (%) 65.6 68.5 77.4

(a) )c()b(

Max. 1.6Max. 3.6Max. 3.3

0 90

90

φ1

Φ

Fig. 3. Orientation distribution functions (ODF) of austenite (φ2=45° section of Euler space, notation by Bunge, contour levels 1–2–3): (a) CR5, (b) CR120, (c) HR30 ( : D {4 4 11}

⟨11 11 8⟩, : Goss {1 1 0}⟨0 0 1⟩).

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from the austenitic volume which was more resistant to MIMT owing tothe different grain size distribution or the texture. From the discussionin the preceding paragraph, it is clear that the nearly random texture ofHR30 should be more influential on the distinct flow behavior. This isadditionally supported by the remaining austenitic fraction in CR5, i.e.,14.9%, which was very close to that in CR120.

The above results were reconfirmed from the estimated austeniticfractions by X-ray diffraction methods [30,31] as shown in Table 5. Thefraction was calculated following the empirical equation which wassuggested to take account of moderate textures [31,32] and producedthe best consistency in the varying textures of this study.

VI

=1. 4

+1. 4γ

I I

αI I

+2

(211)+2

γ γ

γ γ

(220) (311)

(220) (311)(1)

where I is the integrated intensity of the specified reflections, i.e.,(211), (220) and (311) by the specified phases, i.e., α and γ. Theestimated fractions in Table 5 generally exhibit good agreement withthe corresponding ones in Table 1 and those in the above paragraph.More pronounced transformation in CR5 and CR120 is clearly seen. Onthe other hand, a few specimens, such as CR120 before the tensile test,exhibited considerable differences between the results by XRD andEBSD. The larger value by XRD was caused by a smaller Iα (211) due to

the different ferritic texture from CR5 in spite of a nearly identicalaustenitic texture. In the cases of considerable texture development asin the current study, it is known that general XRD-based methods areinaccurate in the measurement of phase volume fractions in spite oftheir superior representativeness from larger sampling volumes [31].Thus, for this study, we preferred EBSD which directly identifies thephases of individual grains. As stated in the experimental section, alarge sampling area of 0.8–3.2 mm2 which included at least 10,000grains was scanned to complement the representativeness of EBSD. Inthis study, the results by EBSD should have sufficient reliability as wellas much better consistency in the various textures.

4.2. Susceptibility to MIMT as a function of austenitic orientation ortexture

The progress of MIMT depends on the externally applied stressstate. This would be well represented in the comparison of uniaxialtension and compression applied on metastable austenite. It isgenerally known that tensile strain is more effective on the accelerationof MIMT than the compressive one [33,34]. On the other hand, therehave been several reports on the dependency of MIMT on austeniticgrain orientation [35–40], from which the dependency on texture couldbe speculated. The existence of a preference to a specific stress state,orientation or texture indicates that the driving force for MIMT shouldhave definite anisotropy. In what follows, an analytical method toevaluate the driving force is briefly presented as a measure of the

0.5.xaM5.3.xaM

(a) )c()b(

Max. 5.6

Fig. 4. ODFs of ferrite (contour levels 1–2–3–4–5): (a) CR5, (b) CR120, (c) HR30 ( : RC {0 0 1}⟨1 1 0⟩, ---: RD//⟨1 1 0⟩, ——: ND//⟨1 1 1⟩, ·····: ND//⟨0 0 1⟩).

Table 4Major texture components found in the current work.

Phase Component Euler angle (φ1, Φ, φ2) (deg.) Miller index {h k l }⟨u v w⟩

γ D (90, 27, 45) {4 4 11}⟨11 11 8⟩Goss (90, 90, 45) {1 1 0}⟨0 0 1⟩

α RC (0, 0, 45), (90, 0, 45) {0 0 1}⟨1 1 0⟩ND//⟨111⟩ (φ1, 55, 45) {1 1 1}⟨u v w⟩-fiberRD//⟨110⟩ (0, Φ, 45) {h k l }⟨1 1 0⟩-fiberND//⟨001⟩ (φ1, 0, 45) {0 0 1}⟨u v w⟩-fiber

ND

RD50 μμm 50 μm

(b)(a)

γα

Fig. 5. Microstructures in uniformly elongated volumes after the tensile tests to fracture (overlay of phase and band contrast maps): (a) CR120, (b) HR30 (remained austenite fraction:16.6% and 29.1% for (a) and (b) respectively).

Table 5Austenitic phase fractions (%) measured by XRD before and after the tensile test.

CR5 CR120 HR30

Before 57.5 70.0 58.4After 16.5 17.8 31.2

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susceptibility to MIMT. The influence of orientation or texture will alsobe explained. Thus, this discussion will be confined to the comparisonof CR120 and HR30 which have different austenitic textures.

With identical austenitic fractions and final annealing tempera-tures, the chemical part of the driving force should be the same inCR120 and HR30. Only consideration of the mechanical part, i.e., themechanical interaction energy (U) between the imposed stress state (σ)and the transformation strain (εtr) [34], should suffice.

U σ ε= :i tr i, (2)

In the above Eq. (1), the superscript i corresponds to the i-thvariant of martensite. A positive value of U means a constructiveinteraction between σ and εtr, which increases the susceptibility toMIMT through the formation of the corresponding martensitic variant.

εtr is obtained from the shape deformation (P1) by transformation,

P RBP=1 (3)

R, B and P respectively correspond to the rigid body rotation, Baindistortion and the lattice invariant shear from the Wechsler-Liberman-Read (W-L-R) crystallographic theory for martensitic transformation[34,41,42]. With the lattice parameters of austenite and martensite andan assumption on their orientation relationship, R, B, P and conse-quently P1 can be obtained [34,43]. In this study, the most conven-tional Kurdjumov-Sachs (K-S) orientation relationship [44] was as-sumed to obtain them [34,43], although experimental ones usuallyexhibit small deviations from it because of their irrational nature[42,45,46]. The necessary lattice parameters were calculated based onthe empirical equations by Cheng et al. [47] who gave relationshipsbetween the lattice parameters and the concentrations of C and N. Theeffect of other elements was ignored, as the influence of the twointerstitial elements should be overwhelming especially in the givenhigh-nitrogen steel.

εtr represented in the martensitic crystal coordinate can beobtained from the above P1 in the same coordinate [34,43].

ε P P I= 12

[ − ]mtr

mT

m( ) 1( ) 1( ) (4)

The subscript (m) indicates that the above tensors are representedin martensitic crystal coordinate. I is the identity matrix and thesuperscript T indicates the transpose of a matrix. εtr in the sample (i.e.reference) coordinate is finally obtained by the rotation of the abovestrain tensor with the orientation of a martensitic variant which can beobtained from the measured austenitic orientation (ga) and thecorresponding variant of the orientation relationship (Δgi) [34,43].

ε g g ε g g=(∆ ) (∆ )tr i ia m

tr ia

T,( ) (5)

In Eq. (1), the orientation-dependency of U can be manifested viathat of σ and εtr. Eq. (4) is the detailed description of the latter effect.This could be shown when a fixed (iso-) stress condition was assumedon all grain orientations, i.e.

⎡

⎣⎢⎢

⎤

⎦⎥⎥

σ σgσ =

= 0 00 0 00 0 0

,for∀ a

11 0

(6)

where σ0 is an arbitrary constant. Fig. 6(a) visualizes the distribution ofthe normalized U (U/σ0, i.e., equivalent to εtr

11) in the orientation spaceunder the iso-stress condition. With more focus on the nucleation ofmartensite, only the maximum values of U among the 24 K-S variantswere considered. The major austenitic texture component of CR120 inFig. 3(b), i.e., D, had the lowest U and should be less favorable forMIMT. Although this result can be supported by a similar one from aprevious study [37], it should contradict the observation in Fig. 5.Therefore, the observed dependency of MIMT on austenitic texture inthe comparison of CR120 and HR30 cannot be elucidated simply by theanisotropy in transformation strain due to texture.

In fact, individual austenitic grains should be subjected to differentstress states depending on their orientations, which also contributes tothe orientation-dependency of MIMT. In an attempt to assess thisfactor, the varying stress state as a function of orientation can becalculated with a kind of simple Taylor-type plasticity model [48–50].In the calculation, a simple stress state was assumed, i.e.,

Fig. 6. Map (φ2=45° section of Euler space) of the orientation effect on mechanically induced martensitic transformation: (a) normalized interaction energy of transformation (U/σ0)

under iso-stress condition, (b) normalized stress (σ11/τ0) and (c) normalized interaction energy (U/τ0) in the case of orientation-dependent stress evolution.

Fig. 7. Effect of texture on mechanically induced martensitic transformation: (a) normalized interaction energy of transformation under iso-stress condition, (b) normalized tensilestress and (c) interaction energy in the case of orientation-dependent stress evolution.

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⎡

⎣⎢⎢

⎤

⎦⎥⎥

σ gσ =

( ) 0 00 0 00 0 0

a11

(7)

where σ11 was a function of austenitic orientation, (ga); thus, it wasdifferent from grain to grain. An identical strain rate along the tensiledirection, i.e., a constant ε11, was imposed on all the austenitic grains(or orientations), and 12 systems of {111}⟨110⟩ slip in the untrans-formed austenitic volume were considered. From these assumptionsand the viscoplastic flow rule [50,51], σ11 was calculated by thefollowing equation.

⎡⎣⎢⎢

⎤⎦⎥⎥σ τ ε

γ m m m m=3

∑ (2 − − )ss s s s n

n

11 011

0 11 11 22 33

1

(8)

τ0 and γ0 are the reference values of resolved shear stress and strain ratefor {111}⟨110⟩ slip systems, respectively, and the latter was set to 1. ε11is the applied strain rate along the tensile direction and was fixed at 1. nis the exponent in the viscoplastic flow rule that corresponds to thereciprocal of rate sensitivity, and the value of 19 was given followingthe routine choice for the deformation of metallic materials at roomtemperature [52,53]. Schmid tensor m of an arbitrary slip system s isdefined by the following equation [50].

m n b n b= 12

( + )ijs

is

js

js

is

(9)

where n and b are respectively the slip plane normal and slip directionof the corresponding slip system s. It should be noted that m isrepresented in the sample coordinate system, so it should be a functionof orientation as n and b rotate according to orientations, which finallygives the orientation-dependency to σ11.

Fig. 6(b) presents σ11 normalized with τ0 according to austeniticorientations after a very small tensile strain of 0.005. It shows analmost exactly opposite tendency to Fig. 6(a) which actually representsthe anisotropy of transformation strain due to orientation. The inter-action energy U could be recalculated under this heterogeneous stresscondition and is presented in Fig. 6(c). The recalculated distribution ofU was substantially different from the initial one in Fig. 6(a). AsFig. 6(c) can be regarded as a multiplication of Fig. 6(a) and (b), it isclear that the orientation-dependency of interaction energy is moreaffected by that of stress. The comparison of Figs. 3(b) and 6(c)suggests that the austenitic texture of CR120 is more favorable forMIMT, which now accords with the observation in Fig. 5.

The above discussions on the orientation-dependency of MIMTsufficiently elucidated the observed texture-dependency. However, forthe completeness of discussion, it is necessary to verify the abovereasoning with respect to the experimental textures of CR120 andHR30. Fig. 7(a) shows the evolution of the sample-averaged (or overall)interaction energy under the iso-stress assumption with increasingstrain and consequent rotation of austenitic texture. As speculatedfrom Fig. 6(a), the overall interaction energy is higher in HR30, whichcontradicts the observation in Fig. 5. The evolution of the overalltensile stress and the interaction energy under the heterogeneous stressassumption are presented in Fig. 7(b) and (c), respectively. They ensurethe lower level of stress and the consequent lower interaction energy ofthe austenitic volume in HR30, i.e., its lower susceptibility to MIMT,which determined its distinct plastic flow behavior in Fig. 2.

The above results make it clear that the orientation- or texture-dependency of MIMT should be primarily caused by the anisotropy instress rather than that in transformation strain. It can be reasoned thatthe latter factor should be much weakened as there are a number ofvariants (24), i.e. choices to maximize the interaction energy for eachaustenitic orientation.

5. Conclusions

In this study, lean DSS specimens with the same composition andphase constitution but different grain sizes and textures were preparedby varying the processing route. The specimens exhibited typical TRIPbehavior due to the metastable austenites and resulted in highstrengths and large elongations in the tensile tests. A strongerdependency of MIMT on austenitic texture than on grain size wasdeduced from the slower progress of MIMT and the distinct flowbehavior of HR30 which had a distinct austenitic texture.

The effect of orientation on the susceptibility to MIMT wasevaluated with the classical theory of martensite crystallography. Inthe two different conditions of stress state, the interaction energy oftransformation, which was the mechanical driving force for MIMT,exhibited different orientation- and texture-dependency.

(1) Under the assumption of uniform stress irrespective of grainorientation, the primary texture component (D) in CR5 andCR120 had near minimum interaction energy. As a result,CR120 exhibited lower overall interaction energy than HR30 ofcomparable grain sizes and near-random austenitic texture, whichwas different from the observed larger susceptibility of the formerto MIMT.

(2) On the other hand, the heterogeneous evolution of stress wascalculated using a simple Taylor-type plasticity model, whichindicated near maximum stress evolution on the D componentand consequently a larger overall stress in CR120. The resultinginteraction energy substantially increased in the D component dueto a large stress and the overall energy of HR30 became lower,which agreed well with the experimental observation.

These analyses made it clear that the orientation- or texture-dependency of MIMT was governed by heterogeneous distribution ofstress according to grain orientation.

Acknowledgement

This work was funded by the Fundamental R &D Program of KoreaInstitute of Materials Science (KIMS) (PNK4680). HNH was supportedby the Engineering Research Center (ERC) program the NationalResearch Foundation of Korea funded by the Ministry of Education,Science and Technology (2015R1A5A1037627).

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