Effect of Carbon Gradient on the Microstructure and Mechanical Properties of Fe-22Mn-C TWIP/TRIP Steels

Morteza Ghasri-Khouzania, Moise Bruhisb and Joseph Robert McDermidc

McMaster Steel Research Centre, McMaster University, Hamilton, Ontario, Canada

aghasrim@mcmaster.ca, bbruhis@mcmaster.ca, cmcdermid@mcmaster.ca

Keywords: High-Mn steels; Twinning; ε-martensite; Deformation

Abstract. High-manganese austenitic steels are promising emerging automotive steels

demonstrating high strength and ductility. The main deformation products observed in these steels

are mechanical twins and ε-martensite, where the dominant deformation products vary quite

strongly with stacking fault energy (SFE), which in turn is a very strong function of the alloy carbon

content. In this research, a Fe-22Mn-0.6C sheet steel was decarburized to achieve a variety of

through-thickness C gradients, thereby varying the dominant deformation products through the

sheet thickness, with the overall objective of producing unique microstructures and mechanical

properties. Microstructural analyses after interrupted tensile testing indicated that the amount of

both mechanical twins and ε-martensite increased with increasing true strain, where the deformation

products changed from mechanical twins at the higher-C core to ε-martensite at the lower-C

surface. The spring-back properties of the C graded steels were also compared with reference to the

effect of differential carbon concentration gradient.

Introduction

High manganese sheet steels offer a combination of high strength and large ductility, which can

attributed to the high strain hardening rate of these steels [[1],[2]], making high-Mn steels strong

candidates for automotive mass reduction applications. Two deformation products are usually

observed during the deformation of high-Mn steels: mechanical twins (giving rise to the so-called

TWinning Induced Plasticity or TWIP effect) and ε-martensite (giving rise to TRansformation

Inducted Plasticity (TRIP)), where the driving force for the formation of each is strongly dependent

on the alloy’s stacking fault energy (SFE) [[3]], which is in turn a strong function of temperature

and steel composition [[4]-[6]]. Allain et al. [[4]] predicted that the γ → ε-martensite transformation

is favored if the SFE is lower than 18 mJ/m2 whereas mechanical twins were predicted to form

when the SFE lay between 12 and 35 mJ/m2. Their calculations for Fe-22Mn-C alloys also

predicted a SFE reduction with lower alloy carbon content. These predictions were subsequently

verified by the experimental work of other authors [[7],[8]], where the deformation products

changed from mechanical twins to mechanical ε-martensite when the carbon content was decreased

from 0.6 wt. % to less than 0.2 wt. %.

The principal objective of this work is to investigate the microstructural and mechanical property

evolution of carbon graded Fe-22Mn-C steels as a function of uniaxial tensile deformation. For this

purpose, carbon gradients from 0 wt.% C at the surface to 0.6 wt.% C at the core were obtain by

decarburization treatments, where the deformation products were predicted to change from

mechanical twins at the core to mechanical and thermal ε-martensite at the surface. These carbon-

graded steels were expected to be attractive candidates for applications which require high bending

resistance due to the higher strain hardening rates previously observed for the lower C content

compositions of these alloys [[2]]. Thus, their bending properties were also evaluated.

Advanced Materials Research Vol. 922 (2014) pp 195-200Online available since 2014/May/07 at www.scientific.net© (2014) Trans Tech Publications, Switzerlanddoi:10.4028/www.scientific.net/AMR.922.195

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Experimental Procedure

The as-received materials comprised cold-rolled Fe-22Mn-0.6C steel sheets with a thickness of

1.5 mm. Three carbon-graded Fe-22Mn-C steels (G1, G2, G3) were obtained using decarburization

heat treatments. Decarburized samples comprised 102 mm 36 mm coupons. All decarburization

experiments were performed in a tube furnace at 1000°C followed by oil quenching. The

decarburization parameters are summarized in Table 1.

Table 1. Decarburization conditions.

Alloy Name Atmosphere Time (h)

G1 CO2 (99.9 %)

4 G2 wet H2 (5.2 %)+Ar (94.8 %) 1 G3 wet H2 (5.2 %)+Ar (94.8 %) 3

Carbon concentration profiles for all steels were obtained using a Jobin-Yvon GDOES

instrument. Successive surface layers were removed by grinding and carbon analyses performed,

allowing concentration profiles versus depth to be constructed.

All mechanical testing utilized a 10 kN Instron 5566 tensile frame. Room temperature uniaxial

tensile testing was conducted at a strain rate of 6.67×10-4

s-1

using sub-size tensile samples per

ASTM E8M [[9]] cut from the decarburized coupons using electric discharge machining (EDM).

Tensile tests were interrupted at defined true strains or went until fracture, where the sample

elongation was monitored using a standard strain gauge with a 25 mm gauge length in all cases. V-

bending tests coupled with deformation mapping using digital image correlation (DIC) were also

carried out, where the DIC was used to follow the deformation and bending angle as a function of

punch displacement. V-bend samples (12.7 mm 38.1 mm) were also cut from the decarburized

coupons using EDM.

Electron Back-Scattered Diffraction (EBSD) was conducted on a JEOL JSM-7000F FEG-SEM

to monitor the microstructural evolution across the samples cross-section as a function of applied

true strain. For all analyses, a working distance of 18 mm and an acceleration voltage of 20 keV

were used and the sample was tilted 70° towards the detector. HKL software was used to operate

and control the EBSD analyses, acquire and collect data and as conduct analyses. EBSD sample

cross sections were initially prepared using standard metallographic procedure followed by Ar ion

milling via a JEOL IB-09010 Cross-Section Polisher.

Results and Discussion

Carbon Profile: Fig. 1 depicts carbon concentration profiles from the surface to the core for all

the graded alloys. For the G1 steel, the carbon concentration at the surface was higher than that of

the G2 and G3 steels, consistent with wet H2 being a stronger decarburizing agent than CO2. It can

also be seen that the carbon concentration gradient of the G2 steel was steeper than the G1 and G3

steels because its decarburization time was shorter (1 h). In the case of the G3 alloy, it should be

noted that the carbon content at the core was significantly lower (approximately 0.35 wt.%) than

that observed for either the G1 or G2 alloys, consistent with the longer decarburizing time

employed with wet H2 as the process gas. In the case of the G1 and G2 alloys, it was expected that

the as-decarburized microstructure would exhibit ε-martensite deformation products at the surface

and would mechanically twin at the core whereas the G3 allow would form ε-martensite at the

surface and would have mixed ε-martensite/twin deformation products at the core [[4],[5],[7],[8]].

Tensile Properties: Typical uniaxial true stress-true strain (σ-ε) plots for all graded alloys are

presented in Fig. 2 and their tensile properties summarized in Table 2. All tensile curves exhibited

continuous yielding. The flow curve for G1 and G2 steels showed the three serration types which

have been attributed to the Portevin-LeChatlier (PLC) or dynamic strain aging (DSA) effect

[[10],[11]]. This indicates that the G1 and G2 steels had sufficient carbon at their core to promote

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the PLC effect. Moreover, these flow curves are consistent with those of a homogenous 0.4 wt.% C

or greater 22 wt.% Mn TWIP steel as determined by several authors [[7],8]. This implies that the

tensile response of the G1 and G2 alloys was dominated by the formation of mechanical twins. For

the G3 steel, however, the representative flow curve did not show any serrations (Fig. 2). This steel

had a higher initial work hardening rate than the other steels but a lower uniform elongation (see

Table 2). As indicated by Liang et al. [[2]] and Yang et al. [[7]], this flow curve is characteristic of

high-Mn steel in which the dominant deformation product was ε-martensite – i.e. a TRIP alloy –

and implies that the flow behavior of this alloy was dominated by ε-martensite formation. These

trends are consistent with the predicted phase make-up of the decarburized alloys, as stated above.

Fig. 1. Carbon concentration profile for the

graded steels. Note that the total sheet

thickness is 1.5 mm or 1500 microns.

Fig. 2. Representative RT true stress-true

strain curves of the graded steels.

Table 2. Summary room temperature tensile properties of the experimental steels.

Alloy Yield Stress (MPa) UTS (MPa) Uniform Elongation (%)

G1 219 936 38.1

G2 239 1022 38.3

G3 206 843 26.3

Microstructural Evolution. Since twins and ε-martensite are microstructurally similar to each

other, they cannot be differentiated by conventional SEM. Thus, EBSD phase maps were taken as a

function of applied strain to determine alloy microstructural evolution during tensile deformation.

For all EBSD maps, it should be noted that FCC austenite is red, HCP ε-martensite blue, FCC

manganese oxide (MnO) green, BCT α׳-martensite white and un-indexed area black. Moreover,

grain boundaries are highlighted in black and twin boundaries in yellow.

Fig. 3 illustrates the cross-sectional microstructures of the as-decarburized and deformed G2 and

G3 alloys. The microstructural observation results for the G1 alloy were similar to those of the G2

alloy and are not presented here. For all as-decarburized alloys, a continuous external MnO layer

was observed arising from Mn selective oxidation during decarburization. A thermal α׳-martensite

layer was present just below the MnO external film in which some internal Mn-oxides were

observed. The Mn content of the α׳-martensite layer was measured by EDS to be approximately 4

wt.%. Thus, α׳-martensite formation can be attributed to Mn depletion of the surface layer [[12]]. As

can be seen in Fig. 3(a) and (e), for the as-decarburized graded steels, the microstructure near the

outer surface comprised a dual phase layer of austenite and thermal ε-martensite where the ε-

martensite volume fraction decreased towards the core. The dual phase layer for the G3 alloy was

thicker than that of the G1 and G2 alloys. This was not surprising considering the differences in

carbon concentration profiles for the alloys, where the carbon concentration was at a lower value for

more of the cross section in the case of the G3 alloy (Fig. 1) and was consistent with the SFE

predictions [[4],[5]]. The core microstructure for all of the graded alloys comprised austenite with

Advanced Materials Research Vol. 922 197

some annealing twins, of which the deeper section of the microstructure seen in Fig. 3(a) was

typical.

For all deformed graded steels, the microstructure near the surface was composed of austenite

and ε-martensite, of which the microstructure seen in Fig. 3(b) is typical, where the volume fraction

of ε-martensite increased significantly with applied strain versus the as-decarburized condition (Fig.

3(a) and (e)). No significant population of deformation twins were detected in the surface layers,

leading to the conclusion that the mechanical γ → ε-martensite transformation dominated in the

lower C surface layers (i.e. 0-250 μm below the surface) of the graded steels. However, among

these steels, the depth into the sheet for which mechanical ε-martensite formation was the dominant

mechanism was highest for the G3 alloy due to the consistently lower C values with depth

associated with this alloy. For all deformed graded steels, the volume fraction of ε-martensite

decreased towards the core while the population of mechanical twins increased. In the depth

interval of 250-500 μm below the surface, typical deformation products were a mixture of

mechanical twins and ε-martensite, as seen in Fig. 3(c). Thus, it can be concluded that the

deformation mode for this layer for all graded steels was a combination of TRIP and TWIP. In the

case of the deformed G1 and G2 alloys, the microstructure at the core contained a high population

of deformation twins with a small amount of ε-martensite, as can be observed in Fig. 3(d). This

implies that the dominant deformation mode for the core of these alloys was the TWIP effect. For

the G3 alloy, however, the deformation mode in the core was a mixture of TRIP and TWIP effects,

as can be seen in Fig. 3(f), due to the lower C values associated with this layer. It also seems that

the twinning density for different grains was not the same because the critical stress for twinning is

highly dependent on grain orientation [[13]].

Fig. 3. EBSD phase maps of experimental alloy cross-sections: (a) as-decarburized G2 (0-250 μm

from surface); (b) G2 deformed to ε = 0.2 (0-250 μm from surface); (c) G2 deformed to ε = 0.2

(250-500 μm from surface); (d) G2 deformed to ε = 0.3 (500-750 μm from surface); (e) as-

decarburized G3 (0-250 μm from surface); (f) G3 deformed to ε = 0.2 (500-750 μm below the

surface).

Bending Properties. In order to calculate the spring-back angle of the experimental alloys, the

bending angle was measured using DIC during bend testing. The bending angle as a function of

time for all graded steels are presented in Fig. 4. At the beginning of the test, when the stroke was

zero, the bending angle was 180°. When the punch stroke increased, the bending angle decreased to

its minimum value which corresponded to the lowest punch position. After the stroke release, the

bending angle increased due to spring-back elastic recovery, determined to be 6.4°, 7.3° and 8.6°

198 THERMEC 2013 Supplement

for the G1, G2 and G3 alloys, respectively. It is generally accepted that an increase in yield stress

for a given Young’s modulus value leads to an increase in spring-back angle [[14]]. Thus, the

similar spring-back angles of the steels can be attributed to their similar values of yield stress (see

Fig. 2).

Fig. 4. Bending angle-time plots for the graded steels.

From the V-bending tests, it was also found that damage occurred at the outer fibre of the G3

steel where the greatest tensile strain existed. This can be attributed to the higher concentration of ε-

martensite at this location, where a large volume fractions of ε-martensite have been shown to be

sites of failure initiation in low-C high-Mn steels [[2]]. The G1 and G2 steels did not exhibit any

damage during the bending test as they contained less ε-martensite at the surface layer.

Conclusions

By investigating carbon graded Fe-22Mn-C steels, the following can be concluded:

1) The as-decarburized steels had a dual-phase microstructure of austenite and thermal ε-

martensite, where the ε-martensite amount decreased from the outer surface to the core. The

trends in the phase make-up for the as-decarburized microstructures were generally consistent

with the existing SFE predictions.

2) The tensile deformation products altered from mechanical ε-martensite at the outer surface to

mechanical twins at the core, consistent with the predictions of available models for

deformation products in high-Mn steels.

3) The tensile behavior of the steels were strongly dependent on the core carbon content and the

volume fraction of the core having a carbon content greater than 0.4 wt.%, where the higher

carbon content core alloy displaying an overall tensile behavior typical of a TWIP alloy and

lower carbon content core alloy displaying behavior typical of a TRIP alloy.

4) The spring-back angles measured from V-bending tests were relatively small for all steels.

Acknowledgements

The authors thank the Natural Sciences and Engineering Research Council of Canada (NSERC)

and the members of the McMaster Steel Research Centre for their financial supports. We also thank

the Materials Technology Laboratory of CANMET for fabrication of the Fe-22Mn-0.6C sheet steel

used in the experiments.

Advanced Materials Research Vol. 922 199

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THERMEC 2013 Supplement 10.4028/www.scientific.net/AMR.922 Effect of Carbon Gradient on the Microstructure and Mechanical Properties of Fe-22Mn-C

TWIP/TRIP Steels 10.4028/www.scientific.net/AMR.922.195

DOI References

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http://dx.doi.org/10.1016/S0749-6419(00)00015-2 [2] X. Liang, J.R. McDermid, O. Bouaziz, X. Wang, J.D. Embury, H.S. Zurob, Microstructural evolution and

strain hardening of Fe-24Mn and Fe-30Mn alloys during tensile deformation, Acta Mater. 57 (2009) 3978-

3988.

http://dx.doi.org/10.1016/j.actamat.2009.05.003 [3] L. Remy, A. Pineau, Twinning and strain-induced FCC→HCP transformation in the Fe-Mn-CrC system,

Mater. Sci. Eng. 28 (1977) 99-107.

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fault energy and the plasticity mechanisms in Fe-Mn-C alloys, Mater. Sci. Eng. A. 387-389 (2004) 158-162.

http://dx.doi.org/10.1016/j.msea.2004.01.059 [5] J. Nakano and P.J. Jacques, Effects of the thermodynamic parameters of the hcp phase on the stacking

fault energy calculations in the Fe-Mn and Fe-Mn-C systems, Calphad. 34 (2010) 167175.

http://dx.doi.org/10.1016/j.calphad.2010.02.001 [6] A. Saeed-Akbari, J. Imlau, U. Prahl, W. Bleck, Derivation and variation in compositiondependent

stacking fault energy maps based on subregular solution model in high-manganese steels, Metall. Mater.

Trans. A. 40 (2009) 3076-3090.

http://dx.doi.org/10.1007/s11661-009-0050-8 [10] L. Chen, H.S. Kim, S.K. Kim, B.C. De Cooman, Localized deformation due to PortevinLeChatelier

effect in 18Mn-0. 6C TWIP austenitic steel, ISIJ Int. 47 (2007) 1804-1812.

http://dx.doi.org/10.2355/isijinternational.47.1804 [11] K. Renard, S. Ryelandt, P.J. Jacques, Characterization of the Portevin-Le Chatelier effect affecting an

austenitic TWIP steel based on digital image correlation, Mater. Sci. Eng. A. 527 (2010) 2969-2977.

http://dx.doi.org/10.1016/j.msea.2010.01.037 [13] I. Gutierrez-Urrutia, S. Zaefferer, D. Raabe, The effect of grain size and grain orientation on deformation

twinning in a Fe-22 wt. % Mn-0. 6 wt. % C TWIP steel, Mater. Sci. Eng. A. 527 (2010) 3552-3560.

http://dx.doi.org/10.1016/j.msea.2010.02.041 [14] K.P. Li, W.P. Carden, R.H. Wagoner, Simulation of springback, Int. J. Mech. Sci. 44 (2002) 103-122.

http://dx.doi.org/10.1016/S0020-7403(01)00083-2