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Acta Materialia 58 (2010) 895–903

Deformation-induced martensitic transformation behaviorin cold-rolled and cold-drawn type 316 stainless steels

Nobuo Nakada a,*, Hidetoshi Ito b, Yoshikazu Matsuoka c, Toshihiro Tsuchiyama a,Setsuo Takaki a

a Department of Materials Science and Engineering, Kyushu University, 744 Motooka, Nishi-ku, Fukuoka 819-0395, Japanb Kyushu University now at Nippon Steel Corporation, 1 Chikkoyawatamachi, Sakai-ku, Sakai 590-8540, Japan

c Kyushu University, 744 Motooka, Nishi-ku, Fukuoka 819-0395, Japan

Received 16 July 2009; received in revised form 1 October 2009; accepted 2 October 2009Available online 31 October 2009

Abstract

We investigate deformation-induced martensitic transformation behavior in cold-rolled and cold-drawn specimens of type 316 stain-less steel. Deformation-induced martensite preferentially nucleates at the twin boundary between the austenite matrix and a deformationtwin. In the cold-rolled specimen, martensite formed at the twin boundary has a Kurdjumov–Sachs (K–S) relationship with both theaustenite matrix and the deformation twin (“double K–S relationship”). In the cold-drawn specimen, two kinds of deformation twinswith different twin planes are typically formed, and therefore deformation-induced martensites are formed where the deformation twinboundaries intersect: martensite thus has an imperfect “triple K–S relationship” with the austenite matrix and the two deformation twins.The complicated crystallographic orientation relationship between austenite and martensite grains strongly restricts the formation ofsome variants of deformation-induced martensites. Because of the difference in number of nucleation sites in the cold-drawn andcold-rolled specimens, martensitic transformation is more enhanced in the former than in the latter.� 2009 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

Keywords: Austenitic stainless steel; Deformation-induced martensite; Deformation twin; Variant restriction

1. Introduction

The mechanical properties of metastable austeniticstainless steel are strongly affected by the behavior of defor-mation-induced martensitic transformation. For example,the work hardening of such stainless steels increases asthe amount of deformation-induced martensite formedduring cold working [1–7]. Since the amount of deforma-tion-induced martensite depends on the stability of austen-ite, the work-hardening ability of austenitic stainless steel isusually controlled by the chemical composition that deter-mines the stability of austenite [8,9]. In addition to thechemical composition, deformation processes such as coldrolling and cold drawing also affect the deformation-

1359-6454/$36.00 � 2009 Acta Materialia Inc. Published by Elsevier Ltd. All

doi:10.1016/j.actamat.2009.10.004

* Corresponding author. Tel./fax: +81 92 802 2960.E-mail address: nakada@zaiko.kyushu-u.ac.jp (N. Nakada).

induced martensitic transformation behavior. In Ref. [10],the authors reported that the amount of deformation-induced martensite formed in a cold-drawn specimen oftype 316 stainless steel is significantly larger than that ina cold-rolled specimen under a given strain. This suggeststhat the deformation microstructure of type 316 stainlesssteel develops differently for different deformation pro-cesses, and this phenomenon causes a change in the nucle-ation sites and growth behavior of deformation-inducedmartensite.

In particular, in austenitic steel with low stacking faultenergy (SFE), deformation twins are expected to form easilyand play an important role in determining the behavior ofsubsequent deformation-induced martensitic transforma-tion. Tsakiris et al. have reported that deformation-inducedmartensite grains nucleate at the deformation twin bound-ary, and there exists a specific crystallographic orientation

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relationship between the austenite matrix, twinned austen-ite and deformation-induced martensite [11]. Therefore, ifthe morphology and crystallographic characteristics ofdeformation twins could be varied by the deformation pro-cess, the behavior of deformation-induced martensite, suchas their frequency of nucleation, site distribution and crys-tallographic orientations (texture), should also be change-able by the deformation processes. In this study, thedifference between deformation-induced martensitic trans-formation behaviors of cold-rolled and cold-drawn type316 austenitic stainless steels is investigated in terms ofthe frequency and number of sites of nucleation. In addi-tion, the restriction on the crystallographic variant of mar-tensites is discussed by taking into account the effect ofdeformation twinning on the transformation behavior.

2. Experimental procedure

Commercial type 316 stainless steel (the chemical com-position of which is given in Table 1) was used as the mate-rial in this study. Specimens in the form of plates (3 or10 mm thick) and wires (3 mm diameter) were cut from acold-rolled ingot of type 316 stainless steel and thenannealed at 1423 K for 1.8 ks; they were subsequently sub-jected to air-cooling to form a recrystallized austeniticstructure. The plate specimens were cold-rolled to reducetheir thickness by up to 98% (equivalent strain er � 3.9),and the wire specimens were cold-drawn to reduce theirsectional area by up to 97% (equivalent strain ed � 3.8)using diamond dies. The equivalent strains given by coldrolling and cold drawing specimens are given by Eqs. (1)and (2), respectively:

er ¼2ffiffiffi

3p lnðt0=tÞ; ð1Þ

ed ¼ 2 lnðd0=dÞ ð2Þ

where t0 and t are the thicknesses of the plate specimens be-fore and after cold rolling, and d0 and d are the diametersof the wire specimens before and after cold drawing,respectively. The microstructure of the specimens was ob-served by optical microscopy, field-emission scanning elec-tron microscopy (FE-SEM) and transmission electronmicroscopy (TEM). The specimens were crystallographi-cally characterized by electron backscatter diffraction(EBSD) in the scanning electron microscope. Data ob-tained by EBSD were analyzed using software programsfor automatic crystal orientation mapping (OIM analysisdeveloped by TSL). Crystallographic orientation mappingwas carried out with a step size of 70 or 100 nm, andobtained data with a confidence index value of greater than0.1 were used for carrying out a detailed crystallographic

Table 1Chemical composition of the steel used in this study.

C Si Mn P Cr Ni Mo

0.03 0.48 0.94 0.02 16.3 10.1 2.1

analysis. Specimen hardness was measured by the Vickershardness test (load 49 N). The volume fraction of martens-ite was estimated from the integral intensity ratio of theX-ray diffraction peaks of austenite to those of martensite.

3. Results and discussion

3.1. Differences in microstructural development and

deformation-induced martensitic transformation behavior

between cold-rolled and cold-drawn specimens

Fig. 1 shows an optical micrograph of the annealedspecimen. It has the microstructure of a recrystallizedaustenitic single structure without d ferrite. The mean aus-tenite grain size is approximately 60 lm. Fig. 2 showschanges in the volume fraction of deformation-inducedmartensite in the cold-rolled and cold-drawn specimensas a function of equivalent strain. It is found that deforma-tion-induced martensite forms more easily by cold drawingthan by cold rolling. For example, an equivalent strain of3.0 is sufficient for the transformation to be completed incold drawing but not in cold rolling. The difference in thedeformation-induced martensitic transformation behaviorsduring cold rolling and cold drawing can be attributed tomicrostructural developments occurring during deforma-tion in these two deformation processes. Fig. 3 shows opti-cal micrographs of the deformation structure in the cold-rolled specimens. The microstructures were observed fromthe transverse direction (TD) and rolling direction (RD). Itis found that the austenite grains elongate horizontally inboth the TD and RD sections and form a pancake-likestructure by cold rolling. Since one-directional deformationbands are observed in each grain, it is mainly the primaryslip system that seems to be activated in the case of coldrolling. Further, microstructural development is alsoobserved in the cold-drawn specimen, as shown in Fig. 4.In this case, microstructures were observed from the cir-cumferential direction (CD) and drawing direction (DD).Elongated austenite grains observed in the CD section

Fig. 1. Optical micrograph of the annealed specimen.

Fig. 2. Changes in the volume fraction of deformation-induced martensiteduring cold rolling and cold drawing.

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are similar to those observed in the TD section. However, afew different deformation bands are observed on other slipsystems, as shown by the white arrows in the figure. In theDD section, two different deformation bands are observedclearly, and frequently intersect within an austenite grain,even in the early stage of cold drawing. As a result, austen-ite grains observed in this section are equiaxial. Fig. 5shows TEM images and diffraction patterns of the cold-rolled (Fig. 5a) and cold-drawn (Fig. 5b) specimens withan equivalent strain of 0.3. The cold-rolled specimen hasa fine lamellar structure with a high dislocation density.The diffraction pattern reveals that this lamellar structure

Fig. 3. Optical micrographs showing the mic

consists of the original austenite matrix and a deformationtwin: the two different diffraction patterns obtained fromthe h1 1 0ic incident direction are symmetric in the{1 1 1}c plane [12]. In contrast, the lamellar structureobserved in the case of the cold-drawn specimen developsalong two directions. The diffraction pattern of this speci-men reveals that two kinds of deformation twins with dif-ferent {1 1 1}c twin planes are formed in the austenitematrix (“double-twin relationship”). TEM characterizationof the cold-rolled and cold-drawn specimens reveals thatthe deformation band observed in the optical micrographsshown in Figs. 3 and 4 corresponds to a bundle of finedeformation twins.

3.2. Characteristics of deformation-induced martensite in

cold-rolled specimen

EBSD carried out in a field-emission scanning electronmicroscope is a powerful technique for performing highlyaccurate crystallographic analysis on a wide scale. There-fore, it is suitable for investigating the crystallographic ori-entation of deformation-induced martensites and theirdistribution in an austenite grain. Fig. 6 shows the crystal-lographic orientation distribution (orientation imagingmap) of the cold-rolled specimen with the equivalent strainof 0.6, obtained by EBDS analysis. This is the critical strainvalue at which deformation-induced martensitic transfor-mation begins in this material (see Fig. 2). Data of defor-mation-induced martensite (body-centered cubic (bcc)phase) are shown in Fig. 6b, where the white lines denotethe {1 1 1}c twin boundary. We confirm that many defor-

rostructural development by cold rolling.

Fig. 4. Optical micrographs showing the microstructural development by cold drawing.

Fig. 5. TEM images and diffraction patterns obtained from the specimens with cold rolling (a) and cold drawing (b) of the equivalent strain of 0.3.

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Fig. 6. Crystallographic orientation imaging maps of the cold-rolled specimen with the equivalent strain of 0.6.

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mation twins and fine martensite grains are formed in thedeformed austenitic structure. The deformation-inducedmartensite grains tend to be formed in the vicinity of theinitial austenite grain boundary (c GB) rather than at thecenter of the austenite grains. This implies that concentra-tion of stress around grain boundaries during deformationenhances the deformation-induced martensitic transforma-tion. Fig. 7 shows magnified images of the square area Ashown in Fig. 6. Fine deformation-induced martensitegrains are found to be formed along twin boundariesbetween the austenite matrix (c-1) and twinned austenite(c-2). Only three kinds of crystallographic variants ofdeformation-induced martensite are formed in this area:Variant 1: a0-1 (blue), Variant 2: a0-2 (red) and Variant 3:a0-3 (green). The crystallographic orientations of c-1, 2and a0-1, 2, 3 are shown using pole figures in Fig. 8a andb, respectively. Pole figure (a) shows that the (1 1 1)c-1

Fig. 7. Magnified crystallographic ima

and (1 1 1)c-2 planes are parallel to each other; this impliesthat they correspond to the twin plane of the deformationtwin. The three parallel h1 1 0i directions on the (1 1 1) twinplane are also confirmed on the great circle (broken line).In contrast, pole figure (b) shows that the a0-1, 2 and 3grains share the ð011Þa0 habit plane, which is parallel tothe twin plane, and that the close-packed directions of mar-tensite grains, h1 1 1ia0, are parallel to those of c-1 and 2grains, h1 1 0ic. These relations are summarized as follows:

ð011Þa0-1;2;3kð111Þc-1;2

½1 � 11�a0-1k½0 � 11�c-1k½1 � 10�c-2

½1 � 11�a0-2k½11 � 1�a0-3k½�110�c-1k½�101�c-2

The crystallographic orientation relationship is schemat-ically illustrated in Fig. 9 using Thompson’s face-centeredcubic (fcc) tetrahedra and bcc cubes. This figure shows that

ging maps of the area A in Fig. 6.

Fig. 8. Pole figures of c-1, 2 (a) and a0-1, 2, 3 (b) in the cold-rolled specimen with the equivalent strain of 0.6.

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the a0-1, 2 and 3 deformation-induced martensite grainshave a Kurdjumov–Sachs (K–S) relationship with bothc-1 and 2 grains; this relationship is referred to as a “dou-ble K–S relationship”. From a geometrical viewpoint, itcan be strictly satisfied only on a twin boundary and noton a random grain boundary. Thus, we believe that asthe activation energy of nucleation is reduced, martensitictransformation should occur preferentially on deformationtwin boundaries compared to on austenite grain bound-aries. The experimental results shown in Figs. 6 and 7 indi-cate the reasonability of this belief.

However, it has not been elucidated why only threekinds of variants are formed in the transformation, despitethe possibility of formation of six kinds of variants, shownin Fig. 10, that satisfy the double K–S relationship on a

Fig. 9. Schematic illustration showing double K–S relationship.

twin boundary [13]. It is probable that an anisotropic strainfield is generated through cold rolling or deformation twin-ning and that it imposes a restriction on the formation ofsome variants.

3.3. Characteristics of deformation-induced martensite in

cold-drawn specimen

Fig. 11 shows the crystallographic orientation imagingmap of the cold-drawn specimen with an equivalent strainof 0.6. In the c-1 grain on the right side of the c GB, two

Fig. 10. Schematic illustration showing the relationship between austenitematrix and six martensite variants formed on a twin plane under a doubleK–S relationship.

Fig. 11. Crystallographic orientation imaging maps of the cold-drawn specimen with the equivalent strain of 0.6.

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kinds of intersecting deformation twins with different{1 1 1}c twin planes (c-2 and 3) are mainly observed. Thedistribution of deformation-induced martensite (Fig. 11b)indicates that martensite grains are formed preferentiallyat the intersection points of these twins which have differ-ent directions. It is also confirmed that martensite grainsare similarly formed at many other twin intersections awayfrom the austenite grain boundary. This result suggests thatwhen deformation twinning occurs in multiple directions, anumber of sites over an entire grain are available for thenucleation of deformation-induced martensite, and thisenhances the martensitic transformation in cold drawingas compared to that in cold rolling.

It is interesting to determine whether martensite formedat the twin intersections has a K–S relationship with threeaustenite grains—the original austenite matrix and the two

Fig. 12. Pole figures of c-1, 2, 3 (a) and a0-1, 2 (b) in the c

deformation twins. If this is possible, then twin intersectionscan be considered as the optimum nucleation sites for defor-mation-induced martensite from a crystallographic view-point. As observed in the orientation imaging map shownin Fig. 11b, the deformation-induced martensite grains canbe roughly classified into two kinds of variants: Variant 1:a0-1 (orange) and Variant 2: a0-2 (purple). Fig. 12 showsthe pole figures of c-1, 2, 3 (Fig. 12a) and a0-1, 2 (Fig. 12b).Due to crystal rotation caused during cold drawing, the crys-tallographic orientations of the variants are scattered morediffusively than in the case of cold rolling (see Fig. 8); how-ever, this figure contains sufficient information for analyzingthe crystallographic orientation relationship between theaustenite and martensite grains. A comparison of orienta-tions of c-1, 2 and 3 grains reveals that the c-1 grain has dif-ferent twin relationships with the c-2 and 3 grains. The

old-drawn specimen with the equivalent strain of 0.6.

Fig. 14. {1 1 1}c Standard stereographic projection showing the {1 1 0}a0

poles of six martensite variants formed on a twin plane under a doubleK–S relationship. The center (1 1 1)c is the twin plane corresponding to thehabit plane.

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relationships of the twin planes of the c-1 grain are(1 1 1)c-1 and 2, denoted by A, and (�1 1 1)c-1 and 3, denotedby B, having the common close-packed direction ofh0 1 1ic. Further, as shown in Fig. 12b, there also exists aR3 twin relationship between a0-1 and 2 variants, with acommon close-packed direction that is parallel to that ofc-1, 2, and 3 grains, as follows:

½1�11�a0-1k½1�11�a00-2k½10�1�c-1k½1�10�c-2k½0�1�1�c-3

ð3ÞIt should be noted here that both a0-1 and 2 variants

seem to have two kinds of habit planes (habit planes Aand B), which are parallel to the two twin planes (twinplanes A and B), respectively.

ð10 � 1Þa0-1kð110Þa0-2ktwin plane A

ð110Þa0-1kð0 � 1 � 1Þa0-2ktwin plane Bð4Þ

When both conditions (3) and (4) are satisfied, both a0-1and 2 tend to have a K–S relationship with each of c-1, 2and 3. Thus, such a relation is termed a “triple K–S relation-ship”, and is schematically illustrated in Fig. 13. In otherwords, martensite rather effectively selects the best orienta-tion that fits well to every grain around it, resulting in astrong restriction on the formation of some variants duringmartensitic transformation in the case of cold drawing.

However, in reality, the martensite grain formed at thetwin intersections cannot have an ideal triple K–S relation-ship with all austenite phases. Thus, this relationship devi-ates slightly from the ideal K–S relationship as explainedbelow. Fig. 14 shows a {1 1 1}c standard stereographic pro-jection showing the six K–S variants on the {1 1 0}a0 poles.Here, the center (1 1 1)c is the twin plane that correspondsto the habit plane (1 1 0)a0. As explained in Section 3.2, allsix variants shown in the pole figure have a K–S relation-ship with the twinned austenite. That is, these variants sat-isfy the double K–S relationship. To satisfy the triple K–Srelationship, martensite must have another habit plane that

Fig. 13. Schematic illustration showing a triple K–S relationship.

is parallel to the twin plane of the second twinned austenite.However, it is found that none of the martensite variants V1–V6 (shown in Fig. 14) have another habit plane that is paral-lel to the second twin plane. For example, when (�1 1 �1)c isthe second twin plane (solid triangle in the upper half of thepole figure), none of the {1 1 0}a0 poles coincide with the sec-ond twin plane, as shown in the pole figure. When V1 or V2 isformed as the variant of martensite, the smallest misorienta-tion between the second twin plane and {1 1 0}a0 is 10.57�.Therefore, the experimental results of the triple K–S rela-tionship suggest that such a small misfit strain would beaccommodated by a large number of dislocations introducedduring cold drawing, and this would result in strong restric-tion on the formation of these two variants at the twin inter-sections within an austenite grain.

In this study, we investigated only cold-rolled and cold-drawn specimens; however, we are certain that in speci-mens subjected to other cold-working processes, such asforging, press forming and repeated stress loading (fati-gue), the distribution of deformation twins would vary.In addition, the SFE and austenite grain size are alsoexpected to have a significant influence on the deformationtwinning behavior. As the result, these factors are expectedto cause a change in the distribution, amount and textureof deformation-induced martensite. In order to improvethe mechanical properties of austenitic stainless steels, itis important to consider not only the stability of the austen-ite phase but also the contributions of SFE and grain sizeand the effect of deformation processes.

4. Conclusions

Differences in the deformation-induced martensitictransformation behavior and crystallographic characteris-tics of martensites between cold-rolled and cold-drawn

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type 316 austenitic stainless steel were investigated. Theobtained results are summarized as follows:

(1) Deformation twinning behavior is significantlyaffected by the deformation process. For example,in cold-rolled specimens, mainly one-directionaldeformation twins are formed within a grain, whereasin cold-drawn specimens, two kinds of deformationtwins with different twin planes are formed.

(2) Deformation-induced martensite preferentially nucle-ates at the twin boundary between the austenitematrix and the deformation twin. In addition, thetwin boundary near the austenite grain boundarytends to be a preferential nucleation site in the cold-rolled specimen. However, in the cold-drawn speci-men, deformation-induced martensite is formed atthe intersections of the two kinds of deformationtwins with different directions. Since the intersectionof twin boundaries yields a large number of nucle-ation sites for deformation-induced martensite, mar-tensitic transformation is more enhanced in colddrawing than in cold rolling.

(3) Deformation-induced martensite formed on a twinplane has a K–S relationship with both the austenitematrix and the deformation twin (double K–S rela-tionship). In this study, three kinds of K–S variantsare observed, despite the possibility of formation ofsix variants that satisfy this relationship. This resultmay be attributed to the generation of an anisotropicstrain field because of cold rolling or deformationtwinning.

(4) When deformation-induced martensite nucleates at atwin intersection, it tends to have a K–S relationshipwith all three austenites if both of its habits planes areparallel to two twinned planes (triple K–S relation-ship). However, from a geometric viewpoint, it isimpossible for martensite to have an ideal tripleK–S relationship. The misfit strain produced in thecase of an imperfect triple K–S relationship must beaccommodated by a large number of dislocationsintroduced during cold drawing.

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