Deformation Mechanisms in Austenitic TRIP/TWIPSteel as a Function of Temperature

STEFAN MARTIN, STEFFEN WOLF, ULRICH MARTIN, LUTZ KRUGER,and DAVID RAFAJA

A high-alloy austenitic CrMnNi steel was deformed at temperatures between 213 K and 473 K(�60 �C and 200 �C) and the resulting microstructures were investigated. At low temperatures,the deformation was mainly accompanied by the direct martensitic transformation of c-aus-tenite to a¢-martensite (fcc fi bcc), whereas at ambient temperatures, the transformation via e-martensite (fcc fi hcp fi bcc) was observed in deformation bands. Deformation twinning ofthe austenite became the dominant deformation mechanism at 373 K (100 �C), whereas theconventional dislocation glide represented the prevailing deformation mode at 473 K (200 �C).The change of the deformation mechanisms was attributed to the temperature dependence ofboth the driving force of the martensitic c fi a¢ transformation and the stacking fault energy ofthe austenite. The continuous transition between the e-martensite formation and the twinningcould be explained by different stacking fault arrangements on every second and on eachsuccessive {111} austenite lattice plane, respectively, when the stacking fault energy increased. Acontinuous transition between the transformation-induced plasticity effect and the twinning-induced plasticity effect was observed with increasing deformation temperature. Whereas theformation of a¢-martensite was mainly responsible for increased work hardening, the stackingfault configurations forming e-martensite and twins induced additional elongation during tensiletesting.

DOI: 10.1007/s11661-014-2684-4� The Minerals, Metals & Materials Society and ASM International 2014

I. INTRODUCTION

AUSTENITIC stainless steels are widely used inengineering applications because of their excellent form-ability in combination with high strength and corrosionresistance. In high-alloy austenitic steels, the addition ofchromium, nickel, and manganese stabilizes the high-temperature austenitic phase of iron down to roomtemperature. Depending on the concentration of thealloying elements, the as-produced microstructure con-sists mainly of fcc—c-austenite, although the bcc ferriteor a¢-martensite and hcp—e-martensite may occur.[1]

Furthermore, the austenite often remains in a metasta-ble state and transforms martensitically under mechan-ical loading into the bcc or hcp state. Consequently, themechanical properties are altered tremendously if aphase transformation occurs during plastic deformation.During the c fi a¢ or the c fi e fi a¢ martensitic trans-formation, the transformation-induced plasticity (TRIP)effect is triggered, yielding higher elongation duringtensile testing.[2] By the formation of new interfaces

through the phase transformation, a pronounced work-hardening effect occurs.[3]

The thermodynamic driving force for the a¢-martens-ite formation is related to the difference of the Gibbsfree energy between austenite and a¢-martensite(DGcfia¢). Since the austenite represents the high-tem-perature phase, which is artificially stabilized by alloy-ing, the driving force increases with decreasingtemperature. For different Ni concentrations, DGcfia¢

is plotted as a function of temperature in Figure 1(a).[4]

The graphs exhibit a linear form in which the alloy withthe highest Ni concentration shows a negative DGcfia¢

below approx. 313 K (40 �C). All other alloys with alower Ni content exhibit a higher driving force for thea¢-martensite formation. Furthermore, the martensitictransformation can be enhanced with the aid ofmechanical stresses acting under plastic deformation,although the level of DGcfia¢ is not high enough tocompensate for the interfacial and strain energy duringa¢-martensite nucleation at the specific testing temper-ature.[5] The value of the Md temperature indicates thetemperature at which enough activation energy isprovided by the deformation—in addition to thechemical driving force—to trigger the martensitictransformation.Through the addition of the alloying elements, the

difference between the Gibbs energy (DGcfie) of the fcccrystal lattice and the hcp crystal lattice becomes ratherlow as well.[6,7] This implies that the formation ofstacking faults (SFs) in the fcc austenite lattice, which

STEFAN MARTIN, Postdoc, and DAVID RAFAJA, Full Profes-sor, are with the Institute of Materials Science, TU BergakademieFreiberg, 09599 Freiberg, Germany. Contact e-mail: stefan.martin@iww.tu-freiberg.de STEFFEN WOLF, formerly Postdoc, ULRICHMARTIN, Professor Emeritus, and LUTZ KRUGER, Full Professor,are with the Institute of Materials Engineering, TU BergakademieFreiberg, 09599 Freiberg, Germany.

Manuscript submitted August 8, 2014.Article published online December 3, 2014

METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 47A, JANUARY 2016—49

alters the ABCABC stacking sequence of the {111}lattice planes into the hcp ABAB sequence, requires lessenergy than in pure fcc iron, which is expressed by a lowstacking fault energy (SFE £50 mJ/m2).[8] According tothe approach of Olson and Cohen,[6] the SFE is definedby

cSF ¼ nq DGc!e þ Estr� �

þ 2r nð Þ; ½1�

where n is the number of faulted lattice planes and q isthe density of atoms per lattice plane. Estr represents theelastic strain energy associated with the local fcc fi hcptransformation, and r is the interfacial energy per unitarea between the faulted volume and the non-faultedmatrix.[9] The applicability of this approach was dis-cussed in detail by Geissler et al.[10] If the influence oftemperature is discussed (Figure 1(c)), the main influ-ence is attributed to DGcfie (Figure 1(b)), whereas theother contributions show only a negligible temperaturedependence.[6]

With a low SFE, special deformation mechanisms inthe austenite are activated that involves SF forma-tion,[11–17] namely single-slip e-martensite formation andtwinning. The dissociation width (ds) of partial disloca-tions depends mainly on the shear stress component onthe glide plane szx and the SFE (cSF) according toByun.[18] Assuming an angle of 30 deg between theBurgers vectors of each partial dislocation and the linevector of the perfect dislocation, ds can be calculated asfollows:

dS ¼ 1

8p� 2� 3m

1� m�

Gb2p

cSF � szxbp�2: ½2�

Furthermore, the shear modulus G, Poisson’s ratio m,and the magnitude of the Burgers vector of the partialdislocations bp are included. The smaller the SFE andthe larger the shear stress, the wider the dissociation ofthe partial dislocations.In the case of a low SFE, dislocations dissociate into

Shockley partial dislocations and, thus, the SFs areformed between the partials in the austenite. As crossslip of widely dissociated dislocations is omitted, thedeformation via single-slip is observed for suitablyoriented glide systems. Thus, conventional dislocationglide and the formation of dislocation cells are onlyobserved on glide systems, where the trailing partialexperiences a higher force than the leading partial.[19,20]

Plastic deformation proceeds mainly in narrowly spaceddeformation bands, which consist of SFs, dislocations,or twins of nanometer size.[21] Because of a high densityof SFs in the deformation bands of the austenite,hexagonal areas are formed, which appear as hcp phase(e-martensite) in X-ray diffraction experiments or inTEM and EBSD investigations.[22,23]

Therefore, a continuous phase transformationsequence c fi e fi a¢ is often found in this type ofsteels. Via the c fi e fi a¢ transformation route, thenucleation of the a¢-martensite needs much less activa-tion energy than the c fi a¢ route. This is easilyexplained by considering the involved crystallographicshear operations described in detail by Schumann.[24]

Both, the direct c fi a¢ and the indirect c fi e fi a¢transformations start with shearing of the austenite by1/6a h112i (a is the lattice parameter of the austenite).Through the formation of the hexagonal stackingsequence, this initial shear sequence is automaticallyproduced by the dislocation movement at low SFE,and only the second shear operation needs to beinitiated by elastic strains.[24]

(a)

(b)

(c)

Fig. 1—Thermodynamic driving forces influencing the deformationbehavior of high-alloy austenitic steels as a function of temperature:(a) difference of Gibbs energy for the martensitic transformationsfrom austenite to a¢-martensite for binary iron nickel alloys (datafrom Ref. [4]); (b) difference of Gibbs energy for the martensitictransformations from austenite to e-martensite for CrNi steels (datafrom Ref. [6]); and (c) the stacking fault energy of CrMnNi steels(data from Ref. [63]).

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Deformation twinning is observed at SFEs between20 and 40 mJ/m2. Because the energy barrier for thee-martensite formation becomes unfavorable, twinningis observed instead.[25,26] The mechanism of twin for-mation can be assumed to be similar to that of thee-martensite formation, which manifests a transition fromthe less-ordered SF arrangement during the e-martensiteformation toward the highly ordered stacking faultformation on every consecutive {111}c lattice plane.[27]

Consequently, the ABC stacking order of the austenite isreversed into the R3 twin orientation. In this way, a highshear strain of 0.707 is produced, which results in largestrains until fracture during tensile testing.[28]

Depending on the testing temperature, differentdeformation mechanisms are activated for the investi-gated material with a SFE of ~17.5 ± 1.4 mJ/m2 atroom temperature.[29] The deformation mechanisms forthe CrMnNi steel investigated are correlated with themechanical properties and are summarized in a schema,which involves the deformation temperature and theSFE.

II. EXPERIMENTAL

Due to the high costs of nickel as alloying element, itis intended to substitute nickel by cheaper manganeseinstead.[30] The chemical composition of the investigatedmaterial (Table I) is based on a patent specification formaximizing energy absorption by the TRIP effect.[31]

The steel alloy was cast into sand dies with dimensionsof 200 9 200 9 16 mm by ACTech GmbH, Freiberg,Germany. After machining, the tensile samples with6 mm diameter were heat treated at 1323 K (1050 �C)for 30 minutes and quenched in an argon stream (atroom temperature) for homogenization and in order toremove deformation layers or martensitic phase trans-formations caused by mechanical manufacturing pro-cess from the outer surface. For tensile testing with astrain rate of 0.0004 s�1, a servo-hydraulic testing device(Zwick 1476) equipped with a temperature chamber wasused. After mechanical testing, the samples were axiallydissected and electrolytically polished using a StruersLectroPol-5 with the A3 electrolyte for 3 seconds.

The as-cast material was characterized by a dendriticsolidification associated with chemical segregations anda grain size of 100 to 1000 lm with a high aspect ratio.Although XRD only shows diffraction maxima referringto the fcc austenite, approximately 2 to 3 vol pct of d-ferrite can be found in the dendritic areas, as the high-temperature bcc phase is stabilized down to roomtemperature by the high chromium concentration duringthe primary ferritic solidification. Neighboring areas areenriched by Mn and Ni changing the deformationbehavior locally, which in general depends on the

chemical composition. As the influence of the individualchemical elements is discussed elsewhere by Jahnet al.,[32] the segregation behavior is not the focus ofthe present paper.For qualitative and semi-quantitative phase analysis,

XRD was carried out in Bragg–Brentano geometry(20 £ 2h £ 150 deg) using an RD7 (Seifert/FreibergerPraezisionsmechanik) with CuKa radiation. A graphitemonochromator in front of the scintillation-detectorwas employed to reduce the fluorescence radiation fromiron. Due to the large grain size of the cast material,only a few (though very large) crystallites satisfy thediffraction condition, so that an ideal powder patternwith the expected intensity ratios of the specific hklreflections could not be obtained. With ongoing defor-mation and after the martensitic phase transformation,more reflections appeared, but poor statistics were stillrecognizable. Consequently, a different method wasutilized to quantify the phase fraction of martensitictransformation after deformation. The ferromagneticbcc phase fraction (ferrite+ a¢-martensite) could bedetermined by magnetic measurements, in which themagnetization of the sample in external magnetic fieldsis measured. Under consideration of the actual samplevolume, the magnetic phase fraction can be determinedfrom reference measurements. The benefit of thismethod lies in the integral measurement over the wholevolume of the coarse-grained specimen.The varying deformation mechanisms were investi-

gated by SEM and EBSD. By using Electron Channel-ing Contrast Imaging (ECCI), even individual latticedefects like extended stacking faults, deformation twins,or dislocation tangles can be depicted.[33–37] A LEO 1530Gemini (Carl Zeiss, Oberkochen) operating in highcurrent mode at 20 keV at 5 and 15 mm workingdistance was used for ECCI and EBSD, respectively.Channel 5 software (Oxford HKL), served for theprocessing of the Kikuchi pattern, the phase discrimi-nation, and the orientation determination. The austenitewas described by fcc iron, a¢-martensite by bcc iron, ande-martensite by the hcp high-pressure modification ofiron.

III. RESULTS AND DISCUSSION

A. Microstructure Investigations of the DeformedSamples

The temperature dependence of the different defor-mation mechanisms could be obtained conveniently bycomparing the phase composition before and afterplastic deformation. Although the steel was mainlyaustenitic in the as-produced state, e- and a¢-martensitecould be formed during plastic deformation. Theextent of the martensitic transformation during plastic

Table I. Chemical Composition of the Investigated Alloy (Optical Emission Spectrometry)

Fe Cr Mn Ni Si Al C N

Weight percent bal. 15.5 6.1 6.1 0.9 0.14 0.03 0.03

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deformation at different deformation temperatures isshown in Figure 2, which depicts the XRD data of thesamples deformed until the end of uniform elongation.At 213 K (�60 �C), a¢-martensite was the dominantphase after the plastic deformation, and only a smallamount of austenite remained. With increasing defor-mation temperature, the amounts of a¢- and e-martensitedecreased, which indicated a change in the deformationmechanism during plastic deformation. At temperaturesabove 333 K (60 �C), the austenite was the phasethat accommodated the main part of the plastic defor-mation. Consequently, it could be concluded that adecreasing driving force with increasing deformationtemperature for the formation of a¢-martensite wasactive. The intensity of the corresponding diffractionlines of a¢-martensite was always superimposed by the 2to 3 vol pct of d-ferrite, which was present in the as-caststate. At 473 K (200 �C), the intensity of 111c was notobserved due to the coarse-grained cast state, althoughit should have comprised the highest intensity.

To reveal and highlight the deformation mechanismsin detail, SEM investigations after reaching uniformelongation and in distinct deformation stages wereconducted. In Figure 3, micrographs and EBSD phasemaps are displayed which present the characteristics ofthe specific deformation mechanisms at the individualtemperatures. The SEM image (Figure 3(a)) showsdeformation bands after 10 pct plastic strain at 213 K(�60 �C). Individual stacking faults (the red circle inFigure 3(a)) indicate the development of these bands bytheir continuous growth from stacking faults formed viapartial dislocation separation as a consequence of theapplied stress. At this specific temperature, the individ-ual stacking faults were particularly wide, because theSFE had the lowest value compared to the other testing

temperatures.[38] Consequently, SFs could even be foundspanned on secondary glide planes between the defor-mation bands (arrow), as already observed by Weidneret al.[33] Hardly any dislocation tangles could beobserved along the deformation bands. This indicatedthat due to the low SFE, the critical dissociation stressneeded to generate large stacking faults is low as well. Ifthe density of the stacking faults in the deformationbands is high enough, the deformation bands areassigned to hcp iron (e-martensite).[22,39] The hcplattice provides a suitable nucleation environment fora¢-martensite, as only a single crystallographic shearoperation is necessary for the phase transformation intothe bcc lattice instead of the two-step shear operationfrom the fcc austenite.[24] Inside the deformation bands,many small a¢-martensite nuclei are formed and not onlyat crossings of deformation bands, as assumed by Olsonand Cohen[40] and Sinclair and Hoagland.[41] Thesea¢-martensite crystallites grow until they reach theinterface of other nuclei, forming larger transformationareas. As the macroscopic strain reaches uniformelongation, nearly the whole volume of the tensilesample was transformed into fine-grained a¢-martensite,as seen in Figure 3(f). Only small ‘islands’ of austeniteremain. This could have been due to two reasons. Thefirst one may be a mechanical shielding by surroundinga¢-martensite, which exhibits higher strength and, there-fore, absorbs the load. The second reason may be giventhrough chemical inhomogeneities in terms of higher Niand Mn concentrations, which would lower the drivingforce for a¢-martensite formation. Further, some large,blocky a¢-martensite grains could be found (blackrectangle in Figure 3(f)). This was related to thespontaneous a¢-martensite formation during the ‘elastic’deformation at the beginning of the tensile experiment.This ‘stress-induced’ a¢-martensite formed parallel to theslip planes of the austenite, which is described in detailin References 42, 43. As caused by the subsequentplastic deformation of the whole crystalline aggregate,these straight blocks became deformed and appearedbent, as can be seen in Figure 3(f).At room temperature, the deformation mechanism

was similar to the one described above, although no‘stress-induced’ a¢-martensite was observed. Hence, thea¢-martensite was formed by plastic deformation of theaustenite. In Figure 3(b), it is obvious that most ofthe stacking faults were confined to the deformationbands, and once again the areas beside the deforma-tion bands did not exhibit a high defect density.Deformation bands were identified as e-martensite byTEM or EBSD. Quantitative phase analysis yieldsapprox. 20 vol pct of e-martensite after 20 pct strain.[44]

With ongoing plastic deformation, the deformationbands grew in width and further new deformationbands form alongside because the motion of the (partial)dislocations in the initial deformation band wasobstructed by the a¢-martensite nuclei. This processcontinued until the whole grains were filled withdeformation bands, which transformed successively intoa¢-martensite (Figure 3(g)).[45]

Above room temperature, the transformation behav-ior changed tremendously. The deformation bands with

Fig. 2—XRD Diffraction patterns of the TRIP/TWIP steel afterplastic deformation until uniform elongation at different testing tem-peratures. Below 333 K (60 �C), a pronounced formation of a¢-mar-tensite and e-martensite was observed, whereas at highertemperatures, the TRIP/TWIP steel remained in the fcc austenitephase.

52—VOLUME 47A, JANUARY 2016 METALLURGICAL AND MATERIALS TRANSACTIONS A

a¢-martensite nuclei were still visible after 15 pct ofplastic deformation (Figure 3(c)). The deformationbands originate in the left side of the image and endedin an area where many dislocations were piled up

(red circle). In contrast to the previously describeddeformation temperatures, conventional motion of reg-ular dislocations (or, rather, the movement of partials,where the dissociation width is not large enough to

20 µm

1 µm

2 µm

2 µm

1 µm

2 µm

(a)

5 %

15 %

55 %

30 %

30 %

40 %

55 %

55 %

30 %

10 %10 %

213 K (-60 °C)

293 K (20 °C)

333 K (60 °C)

373 K (100 °C)

473 K (200 °C)

twin bundles

LAGB

(b)

(c)

(d)

(e)

(f)

(g)

(h)

(i)

(j)

Fig. 3—SEM characterization of the deformation mechanism at different strains and testing temperatures. Left column: inverted ECCI; right col-umn: EBSD phase maps. Gray—austenite (band contrast), yellow—e-martensite, blue—a¢-martensite, black—not indexed, red lines—fcc, R3-twinboundaries indicating deformation twins (all EBSD maps f–j share the same scale bar) (Color figure online).

METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 47A, JANUARY 2016—53

prevent cross slip) seemed to play an important role. Asthe critical stress for the ‘infinitely’ wide partial dislo-cation separation was much higher at these temperaturesdue to the increased SFE,[18] much more dislocationmotion was activated before band formation wasobserved. Thus, a lattice curvature could form ‘bent’deformation bands (Figure 3(c)). Defect agglomerationslike dislocation pile-ups acted as barriers for the (partial)dislocation motion, stopping the deformation bands.[46]

It could not be distinguished from the SEM imageswhether the bands contained e-martensite or twins of theaustenite. However, the EBSD measurement clearlyshowed the occurrence of both phenomena after 55 pctof plastic deformation at 333 K (60 �C) in Figure 3(h).The continuous transition between both phenomena insome deformation bands underlines the assumption thatthe stacking fault arrangement decides the character ofthe deformation band. Either a more or less regularstacking on every second {111} austenite lattice of theactive slip plane creates the e-martensite, or a successivestacking of faults on every {111} plane results in thecoherent twin orientation.[6,10]

When the testing temperature was further increased,the predominant deformation mechanism shiftedtoward twinning. In the SEM image (Figure 3(d)),closely spaced twin lamellae are observed. In additionto the twin packages, dislocations and dislocationtangles represent the most pronounced microstructuralfeature. If dislocation tangles are observed, ‘crossslip’ must be an active mechanism.[47] Obviously, theincreased SFE at 373 K (100 �C) enables a recombi-nation of the partial dislocations if the leading partialdislocation is decelerated by the stress fields from theobstacles.[48] Figure 3(d) shows further that duringtwin growth even low-angle grain boundaries (LAGB)could be overcome. This reflects the fact that benttwins can be observed in the deformation microstruc-tures as LAGBs, and lattice curvature from priordislocation slip can be accommodated during twingrowth by partial dislocation movement.[49] As theindividual twin lamellae could not be resolved, thewhole twin bundles were identified as one uniform twinorientation. In Figure 3(i), only the outer interface ofthe twin bundle with the austenitic matrix is depictedby red lines. Further yellow pixels indicate e-martens-ite, which means that there is no strictly periodicarrangement of the stacking faults. So faulted austen-ite, faulted hexagonal e-martensite, twins, and transi-tional states coexist in the deformation bands. Evensome a¢-martensite can be found, but the total amountdoes not exceed 5 vol pct after reaching uniformelongation.

At 473 K (200 �C), the plasticity is mainly driven bydislocation slip. The corresponding dislocations, thedislocation tangles, and their interaction with grainboundaries are depicted in Figure 3(e). Only some twinbundles are detected in the EBSD phase mapping(Figure 3(j)). No a¢-martensite could be detected andall bcc grains contained d-ferrite from the as-cast state.Obviously, the deformation mechanisms as well as thedislocation slip decline above the 473 K (200 �C) testingtemperature.

B. Correlation with the Mechanical Properties

The different deformation mechanisms cause com-pletely different work-hardening behavior and varyingmaximal strains during tensile testing (Figure 4). Thehighest ultimate tensile strength (UTS) value of nearly1500 MPa was observed at 213 K (�60 �C). Tensilestrength decreased with increasing testing temperature(Figure 4(a)). The tensile strain at the end of theuniform elongation exhibited a maximum at 373 K(100 �C), whereas the lowest strength and elongationwere recorded at 473 K (200 �C). The characteristics ofthe work hardening change drastically with the testingtemperatures (Figure 4(b)). The initial decline of thework hardening in all graphs from Figure 4(b) repre-sents the usual behavior of fcc metals.[50] Still, the much

0 0.1 0.2 0.3 0.40

500

1000

1500

true [-]

true

[MPa

]

*

*

*

*

*

* - interrupted test

ε = 0.0004 s-1.

(a)

213 K (-60°C)

293 K (20°C)

333 K (60°C)

373 K(100°C)

473 K (200°C)

0 0.005 0.01 0.0150

0.025

0.05

0.075

0.1

0.125

0.15

( true- 0.02)/G [-]

/G [-

]

(b)

213 K (-60°C)

293 K (20°C)

333 K (60°C)373 K(100°C)473 K (200°C)

Fig. 4—Mechanical properties determined from tensile tests at differ-ent testing temperatures of the CrMnNi-TRIP/TWIP steel: (a) flowcurves, (b) Kocks–Mecking Plot: work-hardening rate plotted vs truestress normalized by the shear modulus. The asterisks mark thedeformations in the samples, which were used for the microstructureanalysis using SEM, EBSD, ECCI, and XRD.

54—VOLUME 47A, JANUARY 2016 METALLURGICAL AND MATERIALS TRANSACTIONS A

higher work hardening at 213 K (�60 �C) could beattributed to the martensitic transformation prior to thedeformation and at the beginning of the deformationprocess. A higher phase fraction of a¢-martensite couldbe found already at the beginning of the mechanicalexperiments, as the testing temperature was situatedbelow the Ms temperature.[32] Therefore, the athermala¢-martensite (£10 vol pct) was formed during coolingof the sample from room temperature down to 213 K(�60 �C). Due to the sigmoidal shape of the stress–strain curves, the work-hardening rates observed at thetesting temperatures of 213 K and 293 K (�60 �C and20 �C) show an increase and a maximum of hardeningduring the plastic deformation, which is known to beattributed to the deformation-induced a¢-martensiteformation.[3] But the yield strength of a¢-martensite isprobably not that high, as the concentration of intersti-tial elements is rather low. Weidner et al.[51] reportedthat the nano-hardness of the a¢-martensite is not muchhigher than the surrounding austenite. The continuouslyevolving a¢-martensite nuclei shorten the mean free pathof the (partial) dislocations in the deformation bands inthe ‘single slip mode’ (compare Figure 3(b)). Themartensitic nuclei create new interfaces that act as hardobstacles to the dislocation motion.[52] With increasingtesting temperature, the work hardening decreases. At293 K and 373 K (60 �C and 100 �C), the expandedregion with a nearly constant value of the strainhardening is dominated by twinning as the deformationmechanism in low SFE alloys.[53] This reduced work-hardening capability (as compared to the a¢-martensiteformation) is due to the fact that the Burgers vector ofthe dislocations moving in the deformation band isparallel to the partial dislocations, which form thetwinned region in the same deformation band.[20] Onlyafter activation of twinning on the secondary glidesystem, a more pronounced hardening effect isobserved.[54,55]

When the data of the tensile tests were complementedby further testing temperatures,[56,57] the dependence ofthe mechanical parameters on the testing temperaturecould be obtained (Figure 5). The yield strength(YS = r0.02, Figure 5(a)) decreases with increasing test-ing temperature as a consequence of the thermalactivation of the dislocation motion.[50] A drop of theyield strength between 293 K and 203 K (20 �C and�70 �C) is associated with stress-induced a¢-martensiteformation during the elastic deformation of the austen-ite. As the phase transformation is accompanied by anincrease of the specific volume (DV � 2.3 pct[22]), adirectional elongation of the tensile sample in thedirection of loading is achieved spontaneously, resultingin a drop in the measured stress.[42] The increase ofYS at temperatures below 213 K (�60 �C) could beexplained by the athermal a¢-martensite formation(~10 vol pct) during cooling of the sample down to thetesting temperature. Consequently, a higher amount ofinterfacial area is created in the large austenitic grains,and the external load is distributed to the volumefractions of austenite and a¢-martensite.

The ultimate tensile strength increased with decreas-ing deformation temperature; the increase was propor-

tional to the amount of a¢-martensite formed between200 K and 573 K (�70 �C and 300 �C) (Figure 5(b)).Remarkably, the UTS increased between 293 K and333 K (20 �C and 60 �C) when the amount ofa¢-martensite increased tremendously (see Figures 3(g)and (h)). The ultimate elongation reached a maximumbetween 213 K and 473 K (�60 �C and 200 �C). TheTRIP effect could only be attributed to the left part ofthe maximum, as a¢-martensite was formed in sufficientamounts to influence the plastic deformation behaviorbelow 373 K (100 �C). The excellent deformationbehavior through the TRIP effect is often dedicated toa delayed fracture due to the pronounced work hard-ening of the severely deformed grains, which dissipatethe strain into neighboring grains.[58]

However, another aspect needs to be taken intoaccount if the ductility of the sample is to be discussed.The transformation of the a¢-martensite proceeds withinthe deformation bands, which exhibit a high SF density.Assuming a perfect hexagonal stacking order producedby the SFs, a shear strain of approx. 0.35 is achievedthat resembles half the amount of twinning (Figure 6).At around 20 pct strain at room temperature, approx.20 vol pct of e-martensite can be quantified usingXRD.[44] Due to the fact that e-martensite is consumedby a¢-martensite formation, a much higher cumulative

(a)

(b)

Fig. 5—Parameters of the CrMnNi-TRIP/TWIP steel obtained fromthe tensile tests performed at different testing temperatures (cf.Fig. 3): (a) YS—Yield strength = r0.02, UTS—Ultimate tensilestrength, UE—Uniform Elongation; (b) ferromagnetic phase fractionafter testing until UE (additional data compared to Fig. 3 takenfrom Refs. [56,57]).

METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 47A, JANUARY 2016—55

amount of e-martensite can be assumed during plasticdeformation. It is proposed that a distinctive contributionof the increased ultimate elongation values in the ‘TRIPregime’ is dedicated to e-martensite formation. Thehighest ultimate elongation at 373 K (100 �C) and theright side of the maximum of the uniform elongationare clearly due to the deformation-induced twinning andthe TWIP effect,[59,60] as only small phase fractions(£5 vol pct) of e-martensite and a¢-martensite are formed.

C. Transition of the Deformation Mechanisms

From the presented microstructural investigationsand from additional interrupted deformation testsperformed at various temperatures (reported in Refer-ences 39, 44, 45), a scheme was developed that includedall the observed deformation mechanisms (see Figure 7).The deformation mechanisms were correlated with themechanical properties and the SFE. The SFE of theCrMnNi TRIP steel investigated was determined to be17.5 ± 1.4 mJ/m2 at room temperature by Rafajaet al.[29] The SFE can be extrapolated to highertemperatures assuming an average increase of the SFEwith temperature of 0.075 mJ/Km2.[61] The basis of alldeformation mechanisms is dislocation movement,

which is altered significantly by the varying SFE atdifferent testing temperatures. At low strains, in gen-eral, and at high deformation temperatures (‡473 K(200 �C)), only dislocation slip is observed. As theformation of large stacking faults depends on the stressthat is applied on the partial dislocations,[18] theseparation width of the partials is too low to producerecognizable SFs at low strains. At higher temperatures,the SFE is too high, such that it can be compensated bythe applied stress to achieve an infinite partial disloca-tion separation. At temperatures below 473 K (200 �C),a pronounced SF formation can be observed. Thearrangement of the SFs is controlled by DGcfie, which isthe main contribution for the calculation of the SFE andexhibits a strong dependence on temperature.[6]

If the energy difference between the fcc and hcpstructure is close to zero or even negative,[6] thee-martensite formation due to the high density of SFsin the deformation bands can be observed. At temper-atures exceeding room temperature, the SFE andDGcfie increase and the separation width of the partialdislocations decreases. The hexagonal stacking sequencebecomes thermodynamically unfavorable. Conse-quently, hcp stacking is omitted and the SFs are orderedmore closely, with every consecutive lattice plane formingtwins in the deformation bands of the austenite.[11] Thenonly the SFs bordering the austenitic twins form an hcpstacking sequence of four lattice planes.[10] Thus, themicrostructural features such as e-martensite and twinsare produced locally by a distinct geometrical arrange-ment of the SFs. In the investigated steel, the term DGcfie

represents the key term which decides whether e-mar-tensite or twins are formed.[10] A specific SFE thresholdfor the onset of different deformation mechanisms cannotbe concluded, as a continuous transition between partic-ular deformation mechanisms as a function of tempera-ture and applied strain was observed. Nevertheless, thereported SFEs between 18 and 40 mJ/m2 were validatedto facilitate the deformation twinning in austeniticCrMnNi steels.[25,26,62]

At temperatures below 333 K (60 �C), a¢-martensite isformed during plastic deformation. With the onset of

1101 21

111

[ 010]1[ 2 0]1 1

0002

AB

CA

BC

AB

AB

AB

} [1 2]161

AB

BC

AC

B

B

31 [1 2]1

[1 2]161

(b)(a)

SF

SF

SF} [1 2]13

1

[1 2]121}

CA 2

1 [1 2]1

32 [1 2]1

65 [1 2]1

SF

SF

SF

SF

SF

Fig. 6—Comparison of the formation of e-martensite and twins by successive arrangement of SFs via the Shockley partial dislocation movement.The model was constructed from EBSD orientation data of the primary slip system in Fig. 2(h). The e-martensite was formed by a regulararrangement of SFs on every second ð�111Þ plane. (a), the twinned orientation (fcc¢) was achieved by shearing the matrix on every ð�111Þ plane.(b) The cumulative amount of shear and the direction for every SF are indicated with respect to the top layer of atoms. Gray balls represent theatomic positions in austenite, yellow balls the atomic positions in e-martensite and purple balls twinned orientation (Color figure online).

Fig. 7—Schematic overview of the activated deformation mecha-nisms as a function of testing temperature in correlation with theSFE and the mechanical properties of the CrMnNi TRIP/TWIPsteel.

56—VOLUME 47A, JANUARY 2016 METALLURGICAL AND MATERIALS TRANSACTIONS A

the transformation into the bcc phase, work hardeningis significantly increased and higher stress levels areobserved. The a¢-martensite evolves primarily after ahigh SF density has formed in the deformation bands,because the e-martensite provides the best nucleationsites. By the local shearing of the austenitic matrixby the SFs, the two-step transformation process fromfcc to bcc is simplified to a single shear operation,lowering the activation energy of the martensitic phasetransformation.[24] By means of plastic deformation,the nucleation sites for the martensitic phase transfor-mation are produced. Through the influence of thelow stacking fault energy, the c fi e fi a¢ transforma-tion is triggered. Below room temperature, the thermo-dynamic driving force for the a¢-martensite forma-tion was increased. Thus, no intermediate e-martensitewas needed, and a direct transformation c fi a¢ wasobserved.

IV. CONCLUSIONS

High-alloy CrMnNi austenitic TRIP/TWIP steel wasdeformed at temperatures between 213 K and 473 K(�60 �C and 200 �C) by quasi-static tensile testing. Thehighest strength was measured at 213 K (�60 �C),whereas the highest elongation was detected at 373 K(100 �C). Microstructural investigations showed a pro-nounced a¢-martensite formation at low temperaturesand a transition via e-martensite to mechanical twinningat higher temperatures. A continuous transition fromthe TRIP to the TWIP effect was observed. Theevolution of the maximum elongation of the TRIP/TWIP steel was shown to be related to the change in thedeformation mechanism. Dislocation motion undercircumstances of low stacking fault energy generatedlarge stacking faults. The SFE controlled the alignmentof the SFs during plastic deformation. The formation ofe-martensite at the temperatures between 213 K and373 K (�60 �C and 100 �C) and the formation of twinsbetween 293 K and 523 K (20 �C and 250 �C) wereobserved. Twinning was observed in the temperaturerange where the SFE was between 17 and 35 mJ/m2. Theformation of twins at higher testing temperatures andhigher SFE values can be explained by a negligiblecontribution of the term DGcfie, when the shearedvolume retains the fcc structure. A scheme was pre-sented assigning temperature and SFE values to theactive deformation mechanisms. From this scheme, theeffect of the changing deformation mechanisms on themechanical properties could be derived.

ACKNOWLEDGMENTS

This work was financially supported by the GermanResearch Foundation (DFG) within the framework ofthe collaborative research center SFB 799 ‘TRIP-matrix composites.’ The authors wish to thank the Insti-tute of Iron and Steel Technology (TU Bergakademie

Freiberg) for the magnetic measurements and the chemi-cal analysis.

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