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Acta Materialia 80 (2014) 327–340

Applying the texture analysis for optimizingthermomechanical treatment of high manganese

twinning-induced plasticity steel

Christian Haase a,⇑, Luis A. Barrales-Mora a, Franz Roters b, Dmitri A. Molodov a,Gunter Gottstein a

a Institute of Physical Metallurgy and Metal Physics, RWTH Aachen University, Aachen 52074, Germanyb Max-Planck-Institut fur Eisenforschung GmbH, Max-Planck-Str. 1, Dusseldorf 40237, Germany

Received 4 July 2014; accepted 30 July 2014Available online 3 September 2014

Abstract

Texture analysis was applied to determine the optimal processing parameters for the thermomechanical treatment of an Fe–23Mn–1.5Al–0.3C twinning-induced plasticity steel. A simple processing route consisting of cold rolling and recovery annealing was used toexplore the possibility of tailoring the mechanical properties of this steel. The thermal stability of mechanically induced twin boundariesprovided high retained yield strength after recovery annealing. In addition, recovery processes facilitated a significantly improved duc-tility compared to the cold-rolled material. It was shown that the analysis of texture evolution during deformation and annealing can beused as an effective tool to optimize cold rolling degree and annealing conditions. A dislocation-based constitutive model was used inorder to validate that the CuT texture component can be used as an indirect indicator for the evolution of the deformation twin density.Furthermore, simulation results identified recovery as the dominating softening mechanism under the applied annealing conditions.� 2014 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

Keywords: TWIP steel; Texture; Microstructure; Recovery; Deformation twinning

1. Introduction

Since the works of Grassel et al. [1,2] and Frommeyeret al. [3], high manganese twinning-induced plasticity(TWIP) steels have gained strong interest of the automobileindustry for wide usage as structural components, andtherefore have moved into the focus of worldwide steelresearch. These face-centered cubic, fully austenitic steelscontain a high amount of manganese (15–30 wt.%) andtypically, additions of C (0.05–1 wt.%), Al (0–3 wt.%)and/or Si (0–3 wt.%) [4]. This alloying concept results ina low stacking fault energy (SFE) in the range between

http://dx.doi.org/10.1016/j.actamat.2014.07.068

1359-6454/� 2014 Acta Materialia Inc. Published by Elsevier Ltd. All rights r

⇑ Corresponding author. Tel.: +49 241 8026877; fax: +49 241 8022301.E-mail address: haase@imm.rwth-aachen.de (C. Haase).

�20 and �50 mJ m�2 at room temperature [5–7]. As a con-sequence of the low SFE, deformation mechanisms such asdeformation twinning and shear banding become activatedduring deformation in addition to dislocation glide, result-ing in high strain hardening rates and thus in high ductilityand high strength with a typical ultimate tensile strengthand elongation to fracture product of more than50,000 MPa% [8–10].

A major shortcoming of TWIP steels in the fully recrys-tallized, coarse-grained state is their relatively low yieldstrength, which is typically in the range between 200 and400 MPa [2,11]. Although a low onset of plastic deforma-tion facilitates energy-effective part shaping, it is detrimen-tal when the material is applied in crash-relevant structuralcomponents (such as the A- or B-pillar in automobiles).

eserved.

328 C. Haase et al. / Acta Materialia 80 (2014) 327–340

Nevertheless, pre-straining with medium or high degrees ofplastic deformation can still be used to enhance the yieldstrength but this process reduces the remaining ductilitybelow the critical value that is required for the final shapeforming. Apart from precipitation hardening due to micro-alloying [12–14] or grain refinement by recrystallization[15–19], a combination of pre-straining in the form of coldrolling and recovery annealing was proven to be a promis-ing method to obtain significantly increased yield strengthalong with appreciable elongation [20–26]. The grain-scalemicrostructure evolution during this processing procedureis shown schematically in Fig. 1. During cold rolling themicrostructure is very effectively refined by the formationof deformation twins, which is often referred to as adynamic Hall–Petch effect [27]. A subsequent recoveryannealing of the deformed sheet at heat treatment temper-atures and times below the onset of primary recrystalliza-tion reduces the dislocation density and reestablishesductility. The deformation twins, which remain thermallystable during recovery annealing [20,21,28,29], act asstrong barriers for dislocation movement in the same wayas grain boundaries and facilitate a high retained yieldstrength [27,30–37].

In contrast to high and medium SFE materials, whichform a Copper (Cu)-type rolling texture during cold roll-ing, low SFE materials, such as austenitic steel or a-brass,are characterized by a Brass-type rolling texture after coldrolling [38–44]. During texture evolution in TWIP steels,pronounced {123}h634i S and {110}h112i brass compo-nents are readily formed due to dislocation glide. Uponfurther deformation, the {11 0}h100i Goss component,the {55 2}h115i copper twin (CuT) component and thec-fiber (h111i//ND) are strengthened [45–48]. It wasalready proposed by Wassermann [49] that in silver andbrass the CuT component was formed as a consequenceof deformation twinning in {112}h111i Cu-oriented grainsdue to the preferable Schmid factor for twinning in thesegrains. This twinning effect causes the transition from theCu-type to the Brass-type rolling texture with decreasingSFE. The development of the CuT component was alsoobserved in high manganese TWIP steels upon deforma-tion [45,46,50]. In contrast to the other deformation texturecomponents Cu, S, Brass and Goss, the CuT texture com-ponent is almost solely formed by deformation twinning,and therefore can be used as an indirect indication for anincrease of the volume fraction of deformation twins. Dur-ing the recovery stage of a subsequent heat treatment, the

Fig. 1. Schematic diagram of the grain-scale microstructu

main deformation texture components are retained due topreservation of the deformed microstructure and slightlystrengthened due to a decrease of the dislocation density[51–53]. After complete recrystallization, a retained rollingtexture with a high degree of randomization was frequentlyobserved in TWIP steels due to oriented nucleation andannealing twinning, respectively [51,54–61].

In our previous study [25], the occurrence of recoveryprocesses during annealing of a 30% cold-rolled Fe–22.5Mn–1.2Al–0.3C TWIP steel was investigated. In thecurrent study, the benefit of recovery annealing of the samematerial on its mechanical properties was analyzed. Theinfluence of initial cold rolling reductions and heat treat-ment on microstructure evolution, yield strength–ductilitycombination and work-hardening capacity was addressed.Particular focus was put on the possibility of utilizing ananalysis of the texture evolution during cold rolling andannealing as a tool to predict the mechanical behavior ofthe material investigated.

2. Applied methods

2.1. Experimental

2.1.1. Material chemistry and processing

The chemical composition of the TWIP steel investi-gated is given in Table 1. The corresponding SFE was cal-culated to be �25 mJ m�2 using a subregular solutionthermodynamic model [62].

The material was melted in a vacuum induction furnacein argon atmosphere, cast into 100 kg ingots and subse-quently homogenization-annealed at 1150 �C for 5 h in amuffle furnace in order to reduce segregation. Afterwards,the initially 140 mm thick ingots were forged at 1150 �C toa height of 55 mm, followed by an additional homogeniza-tion heat treatment at 1150 �C for 5 h. The forged slabswere then hot-rolled at 1150 �C to a thickness of 2.4 mm.The material was then reheated between each of the 25passes. A laboratory rolling mill was used to cold-roll thealloy at room temperature to thickness reductions in therange between 10% and 80%. Finally, the samples weresubjected to isothermal recovery annealing in an air fur-nace at annealing temperatures of 550 �C or 630 �C, whichprecluded recrystallization kinetics. In order to shorten theannealing time, recrystallization annealing was conductedat 700 �C for either 15 min (30% cold-rolled) or 10 min(40% and 50% cold-rolled).

re evolution during the processing procedure applied.

Table 1Chemical composition of the investigated alloy.

Element Fe C Mn Al Si N P

(wt.%) Bal. 0.325 22.46 1.21 0.041 0.015 0.01

1 In Ref. [68] a slightly different formulation is used for the twin volumeevolution.

C. Haase et al. / Acta Materialia 80 (2014) 327–340 329

2.1.2. Specimens and characterization techniquesSpecimens with the dimensions 10 mm � 12 mm (trans-

verse direction (TD) and rolling direction (RD), respec-tively) were cut from the cold-rolled and annealed sheets.The samples were then mechanically ground with 800,1200, 2400 and 4000 SiC grit paper and mechanically pol-ished using a 3 lm and 1 lm diamond suspension. Forscanning electron microscopy (SEM) and electron back-scatter diffraction (EBSD) the RD–ND (ND: normal direc-tion) section was electropolished at room temperature for20 s at 22 V. For texture analysis and hardness testing,the middle layer of the RD–TD section was electropolishedfor 2 min at 22 V. The used electrolyte consisted of 700 mlethanol (C2H5OH), 100 ml butyl glycol (C6H14O2) and78 ml perchloric acid (60%) (HClO4). The same electrolytewas used for preparing the samples for transmission elec-tron microscopy (TEM). In order to reveal the microstruc-ture by SEM, the specimens were additionally etched atroom temperature using a 2% Nital solution (95 mlC2H5OH and 5 ml HNO3).

SEM and EBSD analyses were performed in a LEO1530 field emission gun SEM operated at 20 kV accelerat-ing voltage and a working distance of 10 mm. EBSD map-pings were generated with a step size of 0.28 lm. The HKLChannel 5 software was utilized for data post-processingand removal of wild spikes and non-indexed points, takingat least five neighbor points into account. Furthermore,EBSD mappings were subdivided into subsets includingonly recrystallized (RX), recovered (RC) or deformed(DEF) grains using an algorithm of the MATLAB�-basedMTEX package [63,64]. The internal grain/subgrain mis-orientation was determined using the grain reference orien-tation (GROD-AO) technique, which takes the averagegrain/subgrain orientation as a reference. An internal grainmisorientation threshold of RX < 1.5� < RC < 6� < DEFwas used [65]. Grains containing fewer than 10 data pointswere disregarded.

In order to prepare the TEM samples, the initial speci-mens were ground to a thin layer of �100 lm, from whichdisks 3 mm in diameter were stamped out. The disks werethen electropolished using a double jet Struers Tenupol-5with a voltage of 29 V at 15 �C. The RD–TD sections ofthe final specimens were analyzed in a JEOL JEM 2000FX II analytical TEM operated at 200 kV.

The crystallographic texture was characterized by meansof X-ray pole figure measurements. Three incomplete(0–85�) pole figures, {111}, {200} and {220}, wereacquired at the mid-layer of the sheet thickness on a BrukerD8 Advance diffractometer, equipped with a HI-STARarea detector, operating at 30 kV and 25 mA, using filterediron radiation and polycapillary focusing optics. The

orientation distribution functions (ODFs) were calculatedin MTEX. The volume fractions of the corresponding tex-ture components were calculated using a spread of 15�from their ideal orientation.

The Vickers hardness (ASTM E384-10e2) of the cold-rolled and annealed samples was examined using a Shima-dzu HMV microhardness tester with a load of 1 kg (HV1).Ten indents per sample were performed.

The mechanical properties of the material in deformed,recovered and recrystallized condition were evaluated byuniaxial tensile tests at room temperature and a constantstrain rate of 10�3 s�1 along the previous rolling directionon a screw-driven Zwick 1484 mechanical testing device.Flat bar tension specimens were used with a gauge lengthof 44 mm, gauge width of 12 mm, fillet radius of 20 mmand variable thickness depending on the rolling degree.

2.2. Simulation setup

The model used for simulations is a finite element model(FEM) implementation of the analytical model described inRef. [66]. To achieve this, a number of modifications simi-lar to those described in Ref. [67] were made. First, the dis-location cell structure was neglected. Second, all evolutionequations were rewritten in a per slip/twin system formula-tion. Since the twin volume fraction is in the focus of thiswork, the evolution equations used to calculate the twinvolume fraction are recalled in the following. The equa-tions for the dislocation part of the model are rather stan-dard and can be found in Ref. [68]. 1

The twinning process is split into two parts, namely twinnucleation and twin growth. For the nucleation, the modelof Mahajan and Chin [69] is adopted. It relies on thereaction of two dislocations to form a twin nucleus:a2h01�1i þ a

2h1 0�1i ¼ 3� a

6h11�2i. The twin nucleation rate

( _N b) per twin system b is calculated by multiplying the total

number density of potential twin nuclei per unit time (pbst),

by the probability that a sufficient stress concentration forthe formation of the nucleus exists (ptw), and by the prob-ability that one of those nuclei grows into a twin (pncs):

_Nb ¼ pbstpncsptw ð1Þ

In the FEM implementation ptw and pncs are calculatedanalogous to Ref. [66]. However, as slip systems are differ-entiated, pst can be calculated individually for each twinsystem b based on the slip activity of the slip systemsinvolved:

pbst ¼

_ca1qa2 þ _ca2qa1

L0

ð2Þ

where _ca1, _ca2 are the shear rates on slip system a1=a2,qa1, qa2 are the dislocation densities on slip system a1=a2

and L0 is the length of the sessile partial dislocations

330 C. Haase et al. / Acta Materialia 80 (2014) 327–340

forming the twin nucleus. In order to make the nucleus

grow, a critical stress sc ¼ cSF3bsþ 3Gbs

L0, where cSF is the SFE

and bS is the Burgers vector of the Shockley partial, hasto be overcome [66]. Since energy is always gained duringgrain growth, it is assumed that twins grow instantaneouslyuntil they encounter an obstacle such as a grain boundaryor a twin on a non-coplanar twin system. A new twin isconsidered to be disk-shaped, where the radial dimensionis based on the average twin spacing. The twin volume isthen given by:

V tw ¼p4

et2 ð3Þ

where t is the average twin spacing and e is the average twinwidth.

Finally the twin volume fraction evolution is calculatedby the product of the nucleation rate with the volume thata new twin occupies, and the untwinned volume:

_f tw ¼ ð1� f twÞ _NV tw ð4ÞFurther details such as the incorporation of temperatureare given in Refs. [66,68].

3. Results

3.1. Determination of processing parameters by texture

analysis

3.1.1. Required deformation and heat treatment for

beneficial mechanical properties

In order to achieve a TWIP steel with high yield strengthalong with high ductility, a suitable level of deformation is

Fig. 2. Texture evolution of the investigated Fe–23Mn–1.5Al–

required to introduce a high fraction of deformation twins,and thus to attain an effective reduction of the mean freeglide distance of dislocations. Subsequent recovery anneal-ing offers the possibility of regaining ductility and must becarried out for a period of time that is, on the one hand,long enough to initiate recovery processes and, on the otherhand, short enough to impede recrystallization [25]. Boththe necessary degree of deformation by cold rolling andthe right heat treatment regime can be determined by anal-ysis of texture evolution. In the following we will show theefficiency of texture analysis for choosing the optimal pro-cessing parameters for the aforementioned approach.

3.1.2. Degree of reduction by cold rollingThe texture evolution of the investigated material after

cold rolling in the range between 10% and 80% thicknessreduction is illustrated by u2 = 45� ODF sections inFig. 2. A schematic illustration of the main texture compo-nents observed in the u2 = 45� ODF section and the corre-sponding definitions of these components are given inFig. 2 (top left corner) and Table 2, respectively. Fig. 3depicts the calculated volume fractions of selected texturecomponents. At low rolling degrees of 10–20%, the compo-nents Cu, S, Goss and Brass developed and formed a weakCu-type texture. Increased rolling reduction (30–50%)facilitated a shift of the maximum intensity of the ODFfrom the Brass component into a position between Brassand Goss ({110}h1 15i G/B) along the a-fiber (Fig. 2).Furthermore, a spread from the Goss towards the CuTcomponent along the s-fiber and a weakening of the Cucomponent were observed. As a consequence, fractions of

0.3C steel during cold rolling, ODF sections at u2 = 45�.

Table 2Definition of texture components illustrated in Fig. 2.

Component Symbol Miller indices Euler angles (u1,U,u2) Fiber

Brass (B) {110}h112i (55,90,45) a, bGoss (G) {110}h100i (90,90,45) a, sCube (C) {001}h100i (45,0,45) /E {111}h110i (0/60,55,45) cF {111}h112i (30/90,55,45) cCopper (Cu) {112}h111i (90,35,45) b, sCopperTwin (CuT) {552}h115i (90,74,45) sa-Fiber h110i parallel to NDb-Fiber h110i tilted 60� from ND towards RDs-Fiber h110i parallel TDc-Fiber h111i parallel ND

CuT Copper Goss S Brass E+F Random0

10

20

30

f [vo

l.%]

10% 20% 30% 40% 50% 60% 70% 80%

Fig. 3. Volume fractions of the main texture components developedduring cold rolling.

C. Haase et al. / Acta Materialia 80 (2014) 327–340 331

the Goss and CuT components increased, whereas those ofthe Cu as well as the S components decreased (cf. Fig. 3).As seen in Fig. 2, with further rolling reduction, at60–80% deformation, a weak c-fiber consisting of the{11 1}h011i E and {11 1}h112i F components developed.The increase of the volume fraction of the E + F compo-nents was accompanied by a stagnation of CuT and a fur-ther decrease of the Cu component as well as an increase ofthe volume fractions of the Goss and randomly orientedgrains (cf. Fig. 3).

The relation between microstructure and texture evolu-tion in a similar cold-rolled TWIP steel is described indetail in our previous paper [50]. Most likely due to thesimilar chemical composition and the same SFE value of�25 mJm�2, similar features of the deformed structureand deformation mechanisms, as in the steel studied previ-ously [50], were found to relate to the specific texture com-ponents in the Fe–23Mn–1.5Al–0.3C steel investigated inthe current work. Deformation at low rolling degrees of10–20% was dominated by dislocation glide, and thusresulted in a Cu-type texture. At medium rolling degrees,i.e. 30–50% thickness reduction, the contribution of defor-mation twinning to the accommodation of strain increasedsignificantly. Consequently, the decrease of the volumefraction of the Cu component was accompanied by anincrease of the R3 twin related CuT component. At high

rolling degrees of 60–80%, the c-fiber components, E andF, developed due to successive rotation of twin-matrixlamellae into the rolling plane and severe shear banding.Furthermore, the decrease of the Cu component and thestagnation of the CuT component reflected further defor-mation twinning in Cu-oriented grains and, on the otherhand, elimination of the twin containing grains by shearbands, which prohibited a further increase of the CuT tex-ture component.

Shear bands, as microstructure heterogeneities of highlocalized shear deformation [70], are hardly capable ofaccommodating further strain during deformation, andtherefore are undesirable microstructure constituents instructural components that require uniform elongation.In addition, shear bands act as preferential nucleation sitesfor primary recrystallization due to the high, localizedstored energy in these bands, and thus accelerate the onsetof recrystallization [71,72], which then results in a decreaseof the yield strength. Furthermore, the elimination of twin-matrix lamellae, and thus of the necessary twin boundariesby shear bands [73], is also an undesired effect for theapproach applied in this study. Due to the aforementionedeffects, shear banding should be avoided during cold rollingwhen applying the processing route introduced in thiswork. Since the volume fraction of shear bands is directlyrelated to the fraction of the E and F texture components,samples with the highest fraction of the E and F texturecomponents, namely the 60%, 70% and 80% cold-rolledsamples, were excluded. In addition to the absence of shearbands in the microstructure, a high density of deformationtwins is required in order to attain a high yield strengthafter recovery annealing. Therefore, the deformed speci-mens with the highest twin density, which corresponds tothe largest volume fraction of the CuT component, werechosen, i.e. the 30%, 40% and 50% cold-rolled samples(highlighted in Fig. 3).

3.1.3. Annealing time determined by texture analysis

The texture evolution in the investigated material after30%, 40% and 50% reduction by rolling during annealingis shown by selected ODF sections at u2 = 45� in Fig. 4.Even though the texture intensities and indices (T) are very

Fig. 4. ODF sections at u2 = 45� of the (a) 30% cold-rolled, (b) 40% cold-rolled and (c) 50% cold-rolled material after cold rolling, recovery annealing,partial and complete recrystallization annealing.

CuT Copper Goss S Brass Random0

10

20

30

40

f [vo

l.%]

50% CR 2 min 15 min 30 min 1 h 2 h 8 h 24 h

Fig. 5. Volume fractions of the main texture components of the 50% cold-rolled material during annealing at 550 �C.

332 C. Haase et al. / Acta Materialia 80 (2014) 327–340

low, a clear trend in texture development during annealingis observed, which is consistent with the behavior of similarTWIP steels [51,54]. During the recovery stage, the textureintensity was slightly strengthened, which was accompa-nied by an increase of the texture indices. During further(recrystallization) annealing, the main texture componentswere retained due to oriented nucleation, whereas addi-tional annealing twinning facilitated evolution of a com-plete a-fiber [51] as well as further randomization [56,74](cf. Fig. 4). This behavior can also be seen from the volumefractions of the main texture components during annealing,which is exemplarily shown for the 50% cold-rolled mate-rial annealed at 550 �C (cf. Fig. 5). As a result of disloca-tions’ recovery, the volume fraction of randomly orientedgrains decreased, whereas the fractions of the main

C. Haase et al. / Acta Materialia 80 (2014) 327–340 333

deformation texture components, such as CuT, S andBrass, increased. As indicated in Fig. 4, this was accompa-nied by the slight strengthening of the texture index. Itmust be noted, however, that the measured change in vol-ume fraction of the deformation texture components doesnot correspond to an increased volume fraction of grainswith specific orientations, but is rather a consequence ofless distorted lattice planes due to lower dislocation den-sity, which in turn resulted in less scattering of X-raysand thus in a higher detected intensity. With progressiverecrystallization, the volume fraction of randomly orientedgrains increased and the deformation texture componentsdecreased. Therefore, annealing for 30 min at 550 �C(highlighted in Fig. 5), after which the fraction of theCuT component was maximum and the fractions of themain deformation texture components still remained atthe highest level, was chosen as optimal heat treatmentparameters for the recovery annealing of the 50% cold-rolled samples. The same approach was applied to the30% and 40% cold-rolled material and resulted in recoveryannealing parameters of 630 �C/10 min and 550 �C/1 h,respectively.

In order to control the reliability of the annealingparameters determined by texture analysis, hardness mea-surements after annealing were conducted and the micro-structure prior to and after recovery annealing wascharacterized. The hardness development of the 30%,40% and 50% cold-rolled material with annealing time isshown in Fig. 6. As seen, recovery-annealing times deter-mined by texture analysis for all three specimens (encircleddata points) were just at the beginning of the steep decreaseof hardness, which can be associated with the onset ofprimary recrystallization.

3.2. Microstructure evolution during recovery annealing

After cold rolling to a thickness reduction of 30%, themicrostructure consisted of grains elongated along the roll-ing direction and contained slip lines, deformation bands

0.1 1 10 100 1000 10000 100000

150

200

250

300

350

400

450

Vick

ers

hard

ness

[HV1

]

Time [s]

30% - 630°C 40% - 550°C 50% - 550°C

Fig. 6. Microhardness evolution of the 30%, 40% and 50% cold-rolledmaterial after annealing at 550 �C and 630 �C for various annealing times.

and/or deformation twins of primary, respectively primaryand secondary twin systems (cf. Fig. 7a). In specimens with40% and 50% reduction the grains were further elongatedand the density of the aforementioned microstructuralfeatures increased. After a deformation of 40% first micro-shear bands developed in individual grains (cf. Fig. 7c ande). The microstructures of the cold-rolled material afterrecovery annealing with annealing times determined by tex-ture analysis are illustrated in Fig. 7b, d and f. It was foundthat the deformed grains with their elongated shape and themicrostructural features observed in the cold-rolled statewere still present after recovery annealing. Since the grainrefinement effect due to the mechanically induced twinboundaries is essential for the efficiency of the processingroute applied in this study, their thermal stability duringrecovery annealing was analyzed qualitatively using TEM(Fig. 8). Regardless of the cold rolling degree and anneal-ing temperature applied, the deformation twins introducedduring cold rolling were found to be thermally stable in thenon-recrystallized grains (cf. Fig. 8a–d). Even though pri-mary recrystallization was locally initiated at the annealingtemperatures chosen, the deformation twins were main-tained in the microstructure unless they were consumedby the growth of recrystallized grains into the deformedmatrix, as illustrated in Fig. 8c.

In addition to the thermal stability of deformationtwins, the occurrence of recovery processes during recoveryannealing was also checked using EBSD measurements, asdescribed in detail in Ref. [25]. The EBSD mapping of the40% cold-rolled and recovery-annealed material is shownin Fig. 9 as inverse pole figure (IPF) mapping with orienta-tions parallel to the sheet normal direction (ND). Due tothe varying intragranular misorientations of RX andnon-RX grains, the EBSD data could be broken down intoRX (cf. Fig. 9c) and non-RX grains by applying an intra-granular misorientation threshold [25,65]. Furthermore,the non-RX grains could be further subdivided into RC(Fig. 9a) and DEF (Fig. 9b) grains using the sameapproach, where DEF grains can be considered as slightlyrecovered grains with a high residual dislocation density,whereas RC grains underwent stronger recovery thanDEF grains. In contrast to the typical equiaxed shape ofthe RX grains (cf. Fig. 9c), both RC and DEF grains inFig. 9a and b, respectively, revealed an elongated grainmorphology, inherited from their previous plastic deforma-tion during cold rolling. The corresponding microtexture ofthe RX, RC and DEF grains is shown in Fig. 9d–f. Themicrotexture of the DEF and RC grains was dominatedby the main texture components Goss, Brass and Cu,whereas the RX grains revealed the same components withlower intensity and further formation of widespread orien-tations leading to texture randomization. The grain bound-ary misorientation profiles of the DEF, RC and RX grainsof the deconvoluted EBSD data of the material after 30%,40% and 50% cold rolling and recovery annealing areshown in Fig. 10a–c. The RX grains of the three differentEBSD mappings were characterized by two peaks at 39�

RD

ND(a)

30 µm

30 µm

(f)

30 µm

(e)

30 µm

(d)

30 µm

(c)

30 µm

(b)

twins

SB

Fig. 7. SEM micrographs of the material after: (a) 30% cold rolling, (b) 30% cold rolling + recovery annealing, (c) 40% cold rolling, (d) 40% coldrolling + recovery annealing, (e) 50% cold rolling and (f) 50% cold rolling + recovery annealing.

334 C. Haase et al. / Acta Materialia 80 (2014) 327–340

and 60�, which indicated 38.9�h101i R9 and 60�h11 1i R3CSL boundaries, respectively. In contrast to the low frac-tion of low angle grain boundaries (H < 15�) of the RXgrains, the fraction of these boundaries was significantlyhigher for the DEF and RC grains. Furthermore, the lowangle grain boundary fraction of the DEF and RC grainswas found to increase with increasing rolling degree, andthus indicated a higher residual dislocation density withhigher rolling reduction. The occurrence of recoveryprocesses is further supported by the lower fraction oflow angle grain boundaries in the interior of the RC grainscompared to the DEF grains.

Finally, it is stressed that the capability of the EBSDtechnique to resolve deformation twins with nanoscalewidth and separation distance quantitatively is stronglylimited, which explains the lower fraction of R3 boundariesin the DEF and RC grains compared to the RX grains.Moreover, the deteriorating indexing rate with increasingdeformation level lowered the detected fraction of R3boundaries in the DEF and RC grains significantly (morein DEF than in RC grains), and thus was not representativefor the true density of deformation twins. Furthermore, the

obtained discrepancy between fractions of R3 boundaries inDEF and RC grains can be understood from the fact that inhighly twinned grains due to the additional accommodationof plastic strain by deformation twinning the dislocationdensity remained at a lower level compared to grainscontaining a lower fraction of deformation twins. As aconsequence, the intragranular misorientation caused bydislocations was smaller in grains with high fraction ofdeformation twins (with respective R3 boundaries) and thusa relatively high number of these grains was detected as RCgrains.

3.3. Mechanical properties

The results of uniaxial tensile tests of the material aftercold rolling, recovery annealing and recrystallizationannealing are illustrated in Fig. 11. With increasing rollingreduction the yield strength increased continuously up to1220 MPa after 50% deformation, whereas the elongationdecreased dramatically, as evident from the engineeringstress–strain curve in Fig. 11a. Regardless of the previousrolling degree, recovery annealing resulted in a decreased

1 µm

(a)RD

TD

600 nm

(b)

1 µm

RX

(c)

1 µm

(d)

twin

matrix

(000)

(e)

Fig. 8. TEM bright field micrographs of the material after: (a) 30% cold rolling + recovery annealing, (b, c) 40% cold rolling + recovery annealing, (d)50% cold rolling + recovery annealing; (e) selected area diffraction (SAD) pattern of (d) with [0 1 1]c zone axis.

Fig. 9. IPF mappings of the 40% cold-rolled + recovery-annealed material broken down into: (a) DEF grains, (b) RC grains, (c) RX grains and (d–f) thecorresponding ODF sections at u2 = 45� of (a–c) (levels: 1.0, 2.0, 3.0, 4.0, 5.0, 6.0, 7.0, 8.0).

C. Haase et al. / Acta Materialia 80 (2014) 327–340 335

0.1

0.2

0.3

0.4

0.5

0.1

0.2

0.3

0.4

0 10 20 30 40 50 600.0

0.1

0.2

0.3

0.4

deformed grains recovered grains recrystallized grains

(a)

Rel

ativ

e fre

quen

cy

(b)

Grain boundary misorientation angle [°]

(c)

Fig. 10. Grain boundary misorientation profiles of the DEF, RC and RXgrains after: (a) 30%, (b) 40% and (c) 50% reduction by cold rolling andsubsequent recovery annealing (630 �C/10 min, 550 �C/1 h and 550 �C/30 min, respectively).

0 10 20 30 40 50 60 700

200

400

600

800

1000

1200

1400

Engi

neer

ing

stre

ss, σ

[MPa

]

Engineering strain, e [%]

0.0 0.1 0.2 0.3 0.4 0.50

1000

2000

3000

4000

True

stre

ss, σ

true [

MPa

] /

Wor

k ha

rden

ing

rate

(dσ/

dε) [

MPa

]

True strain, ε

30% CR 30% CR + RC 30% CR + RX 40% CR 40% CR + RC 40% CR + RX 50% CR 50% CR + RC 50% CR + RX

(a)

(b)

Fig. 11. (a) Engineering stress–strain curves and (b) true stress–true straincurves (dotted lines) and work-hardening rate–true strain curves of theinvestigated Fe–23Mn–1.5Al–0.3C steel after various degrees of coldrolling (CR) and subsequent heat treatment. (RC annealing – 630 �C/10 min after 30% CR, 550 �C/1 h (40% CR), 550 �C/30 min (50% CR); RXannealing – 700 �C/15 min after 30% CR, 700 �C/10 min (40% CR),700 �C/10 min (50% CR)).

336 C. Haase et al. / Acta Materialia 80 (2014) 327–340

yield strength along with significantly improved ductility.This effect was most pronounced for the 50% cold-rolledand recovery-annealed material, where the total elongationincreased by a factor of 12 compared to the cold-rolledstate. Moreover, in comparison with the recrystallizedsamples the yield strength after recovery annealingremained at a high level. The 40% and 50% deformedand recovery-annealed materials revealed a yield strengthof 831 MPa and 929 MPa, respectively, and thus raisedthe yield strength by 250% compared to the recrystalliza-tion-annealed states. These trends can also be observedfrom the true stress–true strain curves in Fig. 11b. In addi-tion to the improved ductility, a clearly improved work-hardening capacity of the material after recovery annealingcompared to the cold-rolled samples was also observed (cf.Fig. 11b).

4. Discussion

In order to achieve the desired combination of high yieldstrength and high ductility, the microstructure introduced

by cold rolling and recovery annealing should consist ofa high density of deformation twins, a low fraction of shearbands and a significantly decreased dislocation densitycompared to the cold-rolled state. As reported above, inthe current work the necessary cold rolling and recovery-annealing parameters to achieve this microstructure weredetermined by means of texture analysis. Due to theincreased volume fraction of the c-fiber components(E + F), which are related to the occurrence of shearbands, materials with rolling degrees in excess of 50% werefound to be unsuitable for the applied approach. This wasconfirmed by SEM images of the microstructure after coldrolling (Fig. 7e), where only occasional grain-scale shearbands were observed after 50% deformation. On the otherhand, the increased volume fraction of the CuT componentin specimens deformed up to 50% thickness reductionindicated an increased density of deformation twins. Thisrelationship was confirmed by the calculation of the twin

0.0 0.2 0.4 0.6 0.80

200

400

600

800

1000

1200

1400

1600

True

stre

ss, σ

true [M

Pa]

True strain, ε

Uniaxial compression Uniaxial compression Plane strain compression

Fig. 12. Experimental (line) and simulated (symbols) true stress–truestrain curves of the investigated Fe–23Mn–1.5Al–C steel under uniaxialcompression (triangles) and plane strain compression (squares) boundaryconditions.

C. Haase et al. / Acta Materialia 80 (2014) 327–340 337

density using the dislocation density-based constitutivemodel [66] described in Section 2.2. The good agreementbetween experimental and simulated flow behavior of thematerial during uniaxial compression testing (Fig. 12) con-firms the reliability of the model. The development of thevolume fraction of deformation twins was determined byplane strain compression simulations using the sameparameter set but plane strain boundary conditions. Theresults of these simulations were compared to the evolutionof the CuT texture component, as shown in Fig. 13. At roll-ing degrees below 20% both the calculated twin density inCu-oriented grains and the measured CuT texture compo-nent remained at a low level and obtained a weak contribu-tion of deformation twinning to the accommodation ofplastic strain. At higher cold rolling degrees (P20%) themeasured CuT texture component and the calculated twindensity in Cu-oriented grains were still in good agreementin both the CuT/twin density development (20–40%) and

0 10 20 30 40 50

0

1

2

3

4

5

6

7

8ΔCuT twin fraction (in Cu-oriented grains) twin fraction (in all grains)

ΔCop

perT

win

[vol

.%]/

twin

frac

tion

[vol

.%])

Rolling reduction [%]

Fig. 13. Dependence of cold rolling degree on the evolution of theexperimental volume fraction of the CuT texture component and of thesimulated volume fraction of deformation twins in Cu-oriented/all grains.

saturation (40–50%) regime, and thus revealed the directrelation between the CuT texture component and the den-sity of deformation twins. Due to the formation of the CuTtexture component as a result of deformation twinning inCu-oriented grains, the evolution of the CuT texturecomponent can in principle only be related to twinning inCu-oriented grains. However, the volume fraction of twinsin grains of all crystallographic orientations (Fig. 13)showed a similar slope. Therefore, measurement of the vol-ume fraction of the CuT component allows an assessmentof the overall material behavior.

The reliability of texture analysis as a tool to estimatethe evolution of the twin density was also confirmed bythe change/increase of the yield strength of the cold-rolledsamples. With increasing rolling reductions from 30% to40% and 50% the yield strength increased continuouslyfrom 882 MPa to 1124 MPa, and 1220 MPa, respectively.This was obviously a result of both an increased dislocationdensity and a higher density of deformation twins that pro-moted dynamic grain refinement.

The optimal recovery annealing times after cold rollingwere also determined using texture analysis of the heat-treated specimens. Since TWIP steels are characterized bya slight texture sharpening during recovery, as indicatedby an increased texture index and a higher intensity ofthe main deformation texture components, and a pro-nounced texture randomization during recrystallization,the optimal microstructure was established by extendingthe annealing time to incipient recrystallization, whichwas indicated by an increase of the volume fraction of ran-domly oriented grains and decrease of intensity of the maintexture components after cold rolling. This texture weaken-ing by recrystallization was also confirmed by an analysisof the microtexture of the few recrystallized grains thatappeared after recovery annealing (cf. Fig. 9f). In orderto test the accuracy of the determination of the optimalrecovery annealing time, both SEM imaging and EBSDanalysis were conducted and showed a significant retentionof the morphology of the grains deformed during cold roll-ing. The recrystallized volume fraction after recoveryannealing was found to be less than 10%, which indicatedthat texture analysis during annealing provided a reliableestimate of the optimal annealing conditions. For thisstudy, the time steps chosen were sufficiently discrete inorder to prove the applicability of the analysis methodused. In the case of requiring a more accurate determina-tion of the transition between recovery and onset of recrys-tallization, annealing treatments with finer time incrementscan be performed or in situ XRD can be utilized for themeasurements.

The results of the mechanical tests convincingly demon-strated the efficiency of the processing route consisting ofcold rolling and recovery annealing. The material withthe highest fraction of the CuT texture component aftercold rolling and with the highest fraction of low angle grainboundaries, which indicates a high dislocation density, (cf.Fig. 10a–c), i.e. the 50% cold-rolled specimen, attained the

30 35 40 45 50400

500

600

700

800

900

1000

1100

1200 YS after CR YS after RC Dislocation density after CR Dislocation density after RC

Rolling reduction (%)

Yiel

d st

reng

th (M

Pa)

Rolling reduction [%]

Δσ0.2,RC

Δσ0.2,RC

Δσ0.2,RX

Δσ0.2,RC

Δσ0.2,RX

Δσ0.2,RX

1x1015

2x1015

3x1015

30 40 50

Dis

loca

tion

dens

ity (m

-2)

Fig. 14. Development of 0.2% yield strength and dislocation density aftercold rolling and after cold rolling + recovery annealing. Dr0.2,RX andDr0.2,non-RX indicate the contribution of recrystallization and recovery tothe overall decrease of the 0.2% yield strength.

CR

RC RX600

800

1000

1200 30% 40% 50%

Rp,

0.2

[MPa

]

338 C. Haase et al. / Acta Materialia 80 (2014) 327–340

highest yield strength in the deformed state. However, thetotal elongation of this specimen in tensile tests was verylow. By contrast, recovery annealing improved the ductilityof the cold-rolled samples significantly (cf. Fig. 11a). More-over, the work-hardening capacity after recovery annealingwas also improved (cf. Fig. 11b). The fact that recoveryprocesses really occurred was clearly evidenced by thedeconvoluted EBSD data. The high number of RC grains,which underwent stronger recovery than the less recoveredDEF grains, was characterized by a lower fraction of lowangle grain boundaries (cf. Fig. 10), and therefore con-tained a lower dislocation density as a consequence of therecovery processes. In order to assess the contribution ofrecovery to the regained ductility, this decrease of the dis-location density, q, was calculated from the difference in0.2% yield strength, Dr0:2, between the material in thecold-rolled and recovery-annealed condition. It wasassumed that Dr0:2 can be described as the sum of the con-tributions of both the decrease of q in the RX (Dr0:2;RX)and the non-RX (recovered) grains (Dr0:2;non�RX).

Dr0:2 ¼ f � Dr0:2;RX þ ð1� f Þ � Dr0:2;non�RX ð5Þwhere f describes the recrystallized volume fraction, whichwas obtained from EBSD data. The initial dislocation den-sity after cold rolling, qCR, was estimated by plane straincompression simulations using the aforementioned consti-tutive model. The softening due to the RX and non-RXgrains was calculated by means of the Taylor equation:

Dr0:2;RX ¼ aMGbffiffiffiffiffiffiffiffiqCR

p � ffiffiffiffiffiffiffiffiqRX

p� �ð6Þ

Dr0:2;non�RX ¼ aMGbffiffiffiffiffiffiffiffiqCR

p � ffiffiffiffiffiffiffiffiqRC

p� �ð7Þ

where a � 0.5 is a geometrical constant, M = 3.06 is theTaylor factor,2 G = 60.45 GPa is the shear modulus,b = 2.55 � 10�10 m is the Burgers vector, qRX ¼ 1010 m�2

is the dislocation density in the RX grains after recrystalli-zation [75] and qCR is the dislocation density in the non-RXgrains after recovery annealing. Combining Eqs. 5–7, weobtain:

qRC ¼ffiffiffiffiffiffiffiffiqCR

p � Dr0:2 � f � Dr0:2;RX

ð1� f Þ � aMGb

� �ð8Þ

The evolution of the 0.2% yield strength and the disloca-tion densities after different cold rolling and recoveryannealing treatments are shown in Fig. 14. By consideringthe contributions of both recrystallization and recovery tothe decrease in yield strength one obtained evidence thatrecovery was the dominating softening mechanism. There-fore, the reduced dislocation density (inset in Fig. 14) dueto recovery can with confidence be identified as the maincontribution to the regained ductility.

2 Due to the comparably weak texture of the Fe–23Mn–1.5Al–0.3CTWIP steel investigated, a Taylor factor of M = 3.06 for material withrandom texture was used. Furthermore, M was kept constant for allcalculations since the texture change in the range between 30% and 50%rolling reduction was marginal.

Furthermore, compared to the recrystallization-annealed material a high retained yield strength wasattained after recovery annealing. This effect was promotedby both the significantly higher dislocation density, asdepicted in the misorientation profiles in Fig. 10, and thehigh retained fraction of deformation-induced twinboundaries. Since these twin boundaries were found to bethermally stable up to the onset of recrystallization (cf.Figs. 7 and 8), the fine-grained microstructure waspreserved during recovery annealing and engendered aneffective grain refinement. Compared to the cold-rolledand recovered samples the recrystallization-annealed mate-rial showed a low yield strength, less than 400 MPa, as aresult of the low dislocation density in the RX grains andthe elimination of the deformation twins.

Finally, as illustrated in Fig. 15, the variation of thedeformation degree and the combination of cold rolling

0 10 20 30 40 50 60 70

400

εtotal [%]

Fig. 15. Relationship between total elongation and 0.2% yield strength ofthe material in the CR, RC and RX condition.

C. Haase et al. / Acta Materialia 80 (2014) 327–340 339

and various annealing procedures was found to be aneffective method to produce TWIP steels with a wide rangeof strength–ductility combinations. The density of defor-mation twins and their retention during recovery annealingcan be utilized to process materials with both high yieldstrength and appreciable elongation. The obtainedmechanical properties demonstrate the benefit of theapplied processing route consisting of cold rolling andrecovery annealing to overcome the deficit of either poorductility or low yield strength after cold rolling or/andrecrystallization, respectively.

5. Conclusions

The possibility of using a simple processing routeconsisting of cold rolling and recovery annealing of anFe–23Mn–1.5Al–0.3C TWIP steel for improving mechani-cal properties was investigated. Texture analysis was suc-cessfully utilized for the optimization of the processingparameters. The following results were obtained.

� Texture evolution during both cold rolling andsubsequent heat treatment allowed us to gain a reliableestimate of (i) the required degree of reduction by coldrolling and (ii) the optimal annealing time. (i) Specimenscontaining undesired shear bands were sorted out byinterpreting the evolution of the volume fraction ofc-fiber-oriented grains. The development/increase ofthe density of deformation twins was tracked by analyz-ing the intensity of the CuT texture component. (ii) Thetransition between texture strengthening during recov-ery and texture randomization during primary recrystal-lization was determined as the optimal annealing timefor the processing approach applied.� The optimal processing parameters could be identified

without using additional microscopy techniques. Theobtained texture data proved to be reliable, and thuslends itself as a promising tool for industrial applicationas a non-destructive, online process control method.� A simple combination of cold rolling and recovery

annealing was found to produce TWIP steel with bothhigh yield strength and appreciable ductility. Due tothe thermal stability of nanoscale deformation twinsduring short time annealing the yield strength remainedat a high level. The occurrence of recovery processes wasverified and the contribution of recovery to the regainedductility was proven. Therefore, a significant retentionof mechanically induced twin boundaries allows tailor-ing of the mechanical properties and generates anextended portfolio of mechanical properties to be real-ized by varying cold rolling degrees and annealing tem-peratures and times.� Simulations of the deformation behavior of the TWIP

steel by utilizing the dislocation-based constitutivemodel provided evidence that the CuT texture compo-nent can be used as an indirect indicator for the evolu-tion of the deformation twin density. Furthermore,

simulation results identified recovery as the dominatingsoftening mechanism under the applied annealingconditions.

Acknowledgements

The authors acknowledge gratefully the financialsupport of the Deutsche Forschungsgemeinschaft (DFG)within the Collaborative Research Center (SFB) 761 “Steel –ab initio; quantum mechanics guided design of new Febased materials”. The authors would also like to expresstheir gratitude to Dr. Weiping Hu for his help withTEM experiments. The help of Dr. Su Leen Wong withCP-FEM simulations is also gratefully acknowledged.

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