recrystallisation behaviour of an fe-mn-c-si-al twip
TRANSCRIPT
Recrystallisation behaviour of an Fe-Mn-C-Si-Al TWIP
Lieven Bracke, Nieves Cabañas-Poy
1 Tata Steel R&D, PO Box 10000, 1970CA IJmuiden, The Netherlands
Keywords: TWIP steel, austenite, grain size, recrystallisation
Abstract. The static recrystallisation behaviour of cold rolled and annealed TWinning Induced
Plasticity (TWIP) steels is important for its industrial production. The recrystallisation kinetics have
been determined for an Fe-Mn-C-Si-Al TWIP steel using hardness measurements and
microstructure analysis: it has been shown that recrystallisation progresses rapidly with increased
annealing temperature. Recrystallisation was faster at higher cold reductions, and a smaller final
grain size was observed at lower annealing temperatures. This indicates that the mechanism is
nucleation dominated at lower temperatures; grain growth at higher temperatures appears similar for
all reductions. The recrystallisation results in a crystallographic texture where the main components
of the cold rolling texture are preserved in the final texture after annealing, although some
randomisation was observed.
Introduction
Fully austenitic high Mn steels have been developed recently to meet the automotive industry’s
need for high strength and improved formability [1,2,3]. These enhanced mechanical properties
allow both downgauging and more complex forming operations for automotive parts. Industrial
production of these novel materials requires an extensive knowledge of the annealing response after
cold rolling in terms of recovery, recrystallisation, grain growth and (Fe,Mn)3-carbide precipitation.
The desired microstructure has a fine grain size for increased strength, and contains no carbides in
order to maintain optimum ductility. In the open literature, only a limited amount of data is
available for this class of materials. A lab study on a 60% cold rolled Fe-18Mn-0.6C-1.5Al TWIP
alloy [4] showed that recrystallisation during a 600s annealing cycle starts at 600°C and is
completed by 700°C. Only a very limited softening stage attributed to recovery was found. At
annealing temperatures between 700°C and 840°C, there was only a limited amount of grain growth
due to grain boundary pinning by (Fe,Mn)3-carbides. Above 840°C, the dissolution of these particles
led to appreciably faster grain coarsening. The (Fe,Mn)3-carbide precipitation nose was found at
800°C for an annealing time of 195s. Very fast recrystallisation kinetics have been observed in a
50% cold rolled Fe-22Mn-C-N TWIP alloy [5] using continuous annealing cycle simulations with
holding times up to 120s, again with only a very limited recovery stage. In this material, it was
shown that the deformation texture is retained during recrystallisation. It was argued that this is a
direct consequence of random sampling of recrystallisation nuclei from the deformed matrix,
statistically leading to retention of the crystallographic texture. The high driving force for nucleation
in combination with the absence of a preferred orientation, lead to a near site-saturated nucleation
mechanism. Grain growth was very limited because of direct impingement of growing grains. No
(Fe,Mn)3-carbides were observed in this material, probably due to the relatively short annealing
times employed. Contrary to Fe-Mn based TWIP steels, the annealing response of AISI316L
stainless steel is very sluggish [6, 7], with full recrystallisation of 40% cold rolled samples occurring
only after annealing at 900°C for 7200s. A striking similarity with the Fe-22Mn-C-N TWIP alloy is
that the cold rolled texture is retained after recrystallisation.
This study reports on the annealing response of Fe-Mn-C-Al-Si TWIP steel as a function of the
annealing cycle and the amount of cold reduction.
Materials Science Forum Vols. 715-716 (2012) pp 649-654Online available since 2012/Apr/12 at www.scientific.net© (2012) Trans Tech Publications, Switzerlanddoi:10.4028/www.scientific.net/MSF.715-716.649
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Experimental procedures
The starting material used for this work was a Fe-Mn-C-Si-Al fully austenitic hot rolled strip
with a thickness of 3mm, produced via a semi-industrial route. The stacking fault energy (SFE) of
the material was calculated to be 27mJ/m². The hot rolled microstructure was fully recrystallised
and was free of (Fe,Mn)3-carbides. Cold rolling to different final thicknesses was performed on a
laboratory rolling mill. A number of continuous annealing simulation cycles were performed on
small samples under a protective atmosphere. The soaking temperature was varied between 550 and
900ºC. The heating rate applied was +5°C/s and the soaking time at temperature was 270s. The
cooling rate was -30°C/s to prevent carbide formation during cooling. Microstructure analysis was
performed using a Zeiss Ultra55 FEG-SEM operated at 15kV. Sample preparation consisted of
mechanical polishing down to 1µm, followed by electrolytical polishing. No etching was applied. In
addition to microstructural characterisation, the progress of recrystallisation and grain growth was
followed using Vickers hardness measurements with a load of 5kg.
Results and discussion
The evolution of the hardness during annealing after 50% cold rolling is illustrated in Figure 1
and some corresponding microstructures are shown in Figure 2. Combining these results allows
identification of the different stages during annealing. At temperatures between 550°C and 600°C,
(Fe,Mn)3-carbides start to form, without significantly changing the cold worked structure. The slight
increase of hardness in this temperature interval might be related to the carbide formation, but it
could also be a consequence of experimental scatter. At temperatures above 600°C limited softening
attributed to recovery occurs, while in the temperature interval between 625°C and 650°C,
recrystallisation becomes the main softening mechanism. It is only from this point on that the
mechanical twins induced by cold rolling start to disappear. At a temperature of about 725°C, the
structure is already fully recrystallised, which is followed by grain growth during annealing at
higher temperatures.
200
250
300
350
400
450
500
550
500 600 700 800 900
Temperature (°C)
Hardness (HV)
Carbides
Recovery
Recrystallisation
t = 270s
Grain growth
Figure 1 Vickers hardness (HV5) measurements on samples annealed at different
temperatures, indicating the different stages of the annealing response (annealing time: 270s
for all).
650 Recrystallization and Grain Growth IV
(a) 600°C, 270s (b) 650°C, 270s
(c) 725°C, 270s (d) 900°C, 270s
Figure 2 SEM micrographs of samples annealed at different temperatures (a) (Fe,Mn)3-
carbide precipitation in unrecrystallised matrix, (b) Partially recrystallised structure, (c) Fully
recrystallised structure, (d) Fully recrystallised structure after grain growth (annealing time:
270s for all). The black arrows indicate (Fe,Mn)3-carbides; the white arrows indicate slip lines
which are probably deformation twins.
The softening of the Fe-Mn-C-Si-Al TWIP steel of this study progresses rapidly with increased
annealing temperature, which is similar behaviour to that reported for Fe-18Mn-0.6C-1.5Al [4] and
Fe-22Mn-C-N [5] TWIP steels. The main difference between the different materials is that the
softening starts at a slightly higher temperature for the current Fe-Mn-C-Si-Al TWIP steel. The
sluggish recrystallisation response – related to the formation of annealing twins – as reported for a
cold rolled AISI316 austenitic stainless steel [6,7] has not been observed, even though a large
amount of annealing twins was formed.
The (Fe,Mn)3-carbides dissolved completely in the temperature interval between 800°C and
850°C. It seems that grain growth is enhanced once these carbides have fully dissolved, but this
should be verified in more detail. However, the observations are similar to those in the Fe-18Mn-
0.6C-1.5Al TWIP alloy [4].
The evolution of the crystallographic texture during recrystallisation is shown in Figure 3. As
expected for a material with a low SFE, the cold rolled starting texture is of the Brass type [5, 7, 8].
During recrystallisation, the main texture components are preserved, but the intensity of the texture
weakens considerably from 6.2 times random for the cold rolled condition to 3.1 times random for
the fully recrystallised material. The Fe-22Mn-C-N TWIP steel also showed retention of the texture
components [5], but in that case the intensity of the main texture components was maintained. In the
Fe-Mn-C-Si-Al TWIP steel, the fraction of twin boundaries in the recrystallised condition was about
30% (measured by EBSD), which could explain the randomisation of the texture [9]. No
Materials Science Forum Vols. 715-716 651
comparative data is available for the Fe-22Mn-C-N TWIP steel. Contrary to the observations during
recrystallisation of a 95% cold rolled AISI 316 austenitic stainless steel [7], no new important
texture components emerged in the Fe-Mn-C-Si-Al TWIP steel. This difference is likely to be a
consequence of the much lower amount of cold deformation applied to the former. As such, no
shear bands were present on which any new components might be expected to nucleate.
Figure 3 ϕϕϕϕ2 = 45° and ϕϕϕϕ2 = 65° sections of the ODF for different stages of the recrystallisation
process.
Figure 4 shows that the onset temperature for recrystallisation decreases with increasing amount of
cold reduction. This clearly illustrates the well known fact that an increasing amount of stored
deformation energy increases the driving force for nucleation of new grains. The effect of cold
rolling on the final grain size after recrystallisation annealing at 750°C and 900°C is shown in
Figure 5. The fact that the grain size after recrystallisation at 750°C is reduced with increasing cold
reduction, indicates that more nuclei have formed, and that this higher nucleation density is not
annihilated by grain growth. Note that the grain growth is probably limited by the presence of
carbides, but that does not change the conclusion that the recrystallisation mechanism is nucleation
dominated. For the samples annealed at 900°C, the final grain size is identical regardless of the prior
amount of cold reduction, indicating that grain growth predominantly defines the microstructure at
this annealing temperature.
600
625
650
675
700
725
750
10% 20% 30% 40% 50%
Cold reduction
Onset T of ReX (°C
)
Figure 4 The effect of different levels of cold reduction on the onset temperature for
recrystallisation.
652 Recrystallization and Grain Growth IV
0
2
4
6
8
10
12
10% 20% 30% 40% 50%
Cold reduction
Grain size (µm)
750°C
900°C
growth dominated
nucleation dominated
Figure 5 The effect of cold reduction on the grain size after annealing at 750°C and 900°C.
As mentioned earlier, the ideal final microstructure for this material would be fine-grained for a
high strength, and essentially free of (Fe,Mn)3-carbides to maintain ductility. The results presented
indicate that the smallest grain sizes can only be obtained in the presence of a small fraction of
(Fe,Mn)3-carbides, so a compromise has to be found for the optimised process.
Summary and conclusions
The annealing response of a cold rolled, fully austenitic Fe-Mn-C-Si-Al TWIP steel has been
described in detail.
It has been shown that recrystallisation progresses rapidly in the temperature interval between
625°C and 725°C. Grain growth was limited, which was at least partially due to the presence of
(Fe,Mn)3-carbides. Once these particles had dissolved (above 800ºC), grain growth was more
dominant.
The rolling texture was essentially retained during recrystallisation, although some randomisation
occurred. This could be caused by the formation of annealing twins. The onset of recrystallisation
occurred at lower temperatures for starting materials with a higher degree of cold reduction. The
resulting recrystallised structure was also finer grained, indicating a nucleation dominated
mechanism.
The results indicated that the optimised processing has to bring a compromise between a fine
grain size for strength and a (Fe,Mn)3-carbide free microstructure for ductility reasons.
Acknowledgements
The authors would like to thank Mr. M. Doll (internship from Lyon University), Mr. F. Twisk
and Dr. M. Aarnts for the experimental work carried out at IJmuiden Technology Centre. The
authors gratefully acknowledge Mr. J. Patel and Dr. J. Butler (Swinden Technology Centre) for their
input. HSD TWIP steels are jointly developed by Corus Staal BV and Salzgitter Flachstahl GmbH.
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