Current Opinion in Solid State and Materials Science 8 (2004) 259 265
Transformation-induced plasticity for high strength formable steels
*
P.J. Jacques
D�partement des Sciences des Mat�riaux et des Proc�d�s, Universit� catholique de Louvain, IMAP, Place Sainte Barbe 2,
B-1348 Louvain-la-Neuve, Belgium
Received 31 July 2004; received in revised form 8 September 2004; accepted 14 September 2004
Abstract
Recent advances in the development of high performance steels presenting improved properties of strength and ductility rely on
the TRIP effect, i.e. on the mechanically-induced martensitic transformation of the retained austenite dispersed in a soft ferrite-based
matrix. As a consequence, the stabilisation and retention of austenite at room temperature have become of primary importance,
leading to specifically designed steel grades and thermal or thermomechanical treatments. Particularly, carbon enrichment of the
austenite during intercritical annealing and bainite transformation was found to be very effective in retaining austenite. This meta-
stable austenite then progressively transforms during straining, bringing about a large increase of the work hardening rate. This
increase results from the stress and strain partitioning continuously evolving with the appearance of the hard martensite.
Ó 2004 Elsevier Ltd. All rights reserved.
1. Introduction anism [4 6]. Indeed, this phenomenon, by bringing
about large enhancements of the work-hardening rate,
Over the years, several solutions have been imagined postpones the onset of necking and thus improves the
to improve the mechanical properties of iron and steel in formability.
order to meet the requirements of more and more strin- The TRIP effect has recently regained attention in
gent applications. Indeed, depending on their chemical the case of low alloy steels [7]. The design of new steel
compositions, Fe C based alloys present the exceptional grades and microstructures is indeed motivated by the
feature that the processing route can be adapted to lead necessity for the steel industry to process always better
to various phases that exhibit antagonist mechanical suited high strength structural steels with low produc-
properties ranging from soft and ductile ferrite to ultra tion costs. Responding to an unceasing demand from
high strength martensite. Furthermore, these phases the automotive industry for steels with strength level
can be combined within finely grained microstructures higher than 500 MPa and up to 1000 MPa without sac-
considered as in situ composites. rificing the formability properties, the 1990s have seen
Among the different deformation mechanisms, the the development and characterisation of new formable
martensitic transformation of austenite during mechan- high strength steel grades, the so-called TRIP-assisted
ical loading has been known for quite a long time [1 3]. multiphase steels [8]. As illustrated in Fig. 1, these steels
The acronym �TRIP� (for TRansformation-Induced present complex multiphase microstructures consisting
Plasticity) was proposed to express the efficiency of of a ferritic matrix and a dispersion of multiphase
the martensitic transformation as a deformation mech- grains of bainite, martensite and metastable retained
austenite.
Once the potential improvements brought by these
*
TRIP-aided steels demonstrated, research focussed on
Tel.: [phone]; fax: [phone].
E-mail address: [e-mail] two main axes:
1359-0286/$ - see front matter Ó 2004 Elsevier Ltd. All rights reserved.
doi:10.1016/j.cossms.2004.09.006
260 P.J. Jacques / Current Opinion in Solid State and Materials Science 8 (2004) 259 265
austenite grains dispersed in a soft ferritic matrix. It is
thus required to characterise not only the TRIP effect
but also the composite strengthening effect emerging
from these complex microstructures.
It is proposed hereunder to focus on these two axes,
i.e. to present the current general principles governing
the processing of multiphase microstructures containing
the appropriate austenite and to highlight the main
factors governing their mechanical properties.
2. Processing of TRIP-assisted multiphase steels
In the present case of the TRIP-assisted multiphase
steels, the partial stabilisation of austenite at room tem-
perature is ensured by its carbon content, which is one
of the strongest austenite stabiliser elements. As shown
in Fig. 2, the austenite C enrichment occurs all along
specifically designed thermal or thermomechanical cy-
cles. After cold or hot rolling, the general morphology
of the multiphase microstructure, i.e. the dispersion of
ferrite and austenite grains, results from an intercritical
annealing. The first carbon enrichment of the austenite
accompanies the nucleation and growth of the phases.
Several studies on the formation of the ferrite/austenite
mixture during intercritical annealing have been carried
out in the 1970s and 1980s [13 15] in the case of Dual
Phase steels. However, the maximum austenite C enrich-
ment during intercritical annealing does not prevent the
martensitic transformation on quenching to room tem-
perature. A second stage for further carbon enrichment
is therefore needed.
As shown in Fig. 2, a second hold at an intermediate
temperature combined with particular alloying elements
Fig. 1. SEM (a) and EBSD (b) micrographs illustrating the typical
allow further C enrichment and the stabilisation of aus-
microstructure of the TRIP-assisted multiphase steels ((a) Nital
tenite at room temperature [16,17]. It is well known that
etching; (b) phase map (BCC in grey, FCC in white)).
(i) the design of processing routes bringing about
finely grained microstructures containing the right
amount of austenite with the adapted stability,
(ii) the understanding of the work hardening and
mechanical properties of composite microstructures
exhibiting a TRIP effect.
The TRIP-aided steels actually correspond to a muta-
tion in the way of looking at retained austenite. Indeed,
while large efforts were made to avoid the presence of re-
tained austenite, particularly in welds, for embrittlement
and fracture reasons, recent studies dealt with the max-
imisation of the amount of retained austenite presenting
the right stability [9]. On the other hand, it is also of
primary importance to understand the unique combina-
tions of strength and ductility exhibited by the TRIP-
aided steels. Contrarily to the full TRIP steels developed
Fig. 2. Schematic representation of the thermomechanical treatments
in the 1960s and 1970s [10 12], they indeed present the
applied to hot or cold rolled TRIP-assisted multiphase steels (c:
unique feature that the TRIP effect occurs for small austenite, a: ferrite, a0: martensite, ab: bainite).
P.J. Jacques / Current Opinion in Solid State and Materials Science 8 (2004) 259 265 261
in the course of the bainite transformation, carbon
redistributes from the bainitic ferrite to the surrounding
austenite [18 22], leading to the characteristic bainitic
morphology of ferrite/carbide mixture. Furthermore,
the addition of some alloying elements is known to
promote the formation of carbon-saturated austenite
instead of cementite precipitation. Bainite transforma-
tion has thus become of primary importance in the
processing of the TRIP-aided steels.
Recent work [9] has shown that the bainite transfor-
mation occurring in the case of the TRIP-aided steels
presents some peculiarities. Current issues therefore in-
volve the influence on the austenite retention at room
temperature (i) of the alloying elements [16,17], (ii) of
the austenite grain size [23] and (iii) of the austenite
prior deformation [24].
Beside carbon (from 0.1 to 0.4 wt.%), and manganese
(0.5 to 1.5 wt.%, providing some hardenability), a real
challenge is to find the right alloying element that delays
or suppresses the cementite precipitation during the bai-
nite transformation. It is well known that silicon inhibits
cementite precipitation [21,25]. However, the high sili-
con levels that are needed do not fit well with the indus-
trial practice of galvanised flat products [26]. Aluminium
has also been shown to inhibit effectively the cementite
precipitation and thus to promote the austenite carbon
enrichment [17,27,28]. Some studies dealt with other ele-
ments like Cu [29], P[29,30] or Ni [31]. However, hardly
anything can be found in the literature on the real mech-
anism by which some alloying elements inhibit this
cementite precipitation.
Fig. 3(a) illustrates the evolution of the nature of the
phases constituting the room temperature microstruc-
ture of a typical Si-alloyed TRIP-aided steel as a func-
tion of the progress of the bainite transformation. It is
Fig. 3. (a) Transformation map and austenite C enrichment during the
worth noting that Al- or Al Si alloyed grades exhibit
bainitic hold of a classical TRIP-assisted multiphase steel; (b)
the same transformation behaviour [17]. Depending on
Comparison with the calculated T0 and Ae3 curves of the measured
maximum carbon content of retained austenite of a TRIP-aided steel
the bainitic hold time, various mixtures of bainite, mar-
held at several temperatures. The shaded box corresponds to the range
tensite and retained austenite can be found in the micro-
of carbon content bringing about an optimised mechanical stability of
structure after quenching. The intercritical austenite
the retained austenite.
progressively transforms to bainite, bringing about a
large stabilisation of austenite at room temperature at
the expense of martensite. As also shown in Fig. 3(a), and the maximum carbon content of retained austenite.
this stabilisation is due to the austenite C enrichment The behaviour of these Si- and Al-alloyed steels can thus
operating up to a maximum depending on the tempera- be explained by considering (i) the displacive mechanism
ture at which the bainite transformation takes place. of bainite formation, (ii) the carbon partitioning be-
Some maximum austenite C enrichments measured tween bainitic ferrite and residual austenite and (iii)
when the bainite transformation stops at different tem- the inhibition of cementite precipitation from austenite
peratures for a 1.5Si 1.5Mn grade are given in Fig. [20,22,32].
3(b), together with the T0 and Ae3 curves calculated A shaded box has also been represented on Fig. 3(b).
by the Calphad method. The T0 curve represents the It corresponds to the range of carbon content bringing
upper bound of the austenite carbon content allowing about an optimised mechanical stability of austenite
a composition-invariant transformation of austenite to [33]. Indeed, it is worth remembering that the bainite
ferrite (i.e. where austenite and ferrite of the same com- transformation aims at making the remaining austenite
position present the same chemical free energy). Fig. more stable so that it will progressively transform dur-
3(b) shows a perfect agreement between the T0 curve ing the subsequent forming operations and thus improve
262 P.J. Jacques / Current Opinion in Solid State and Materials Science 8 (2004) 259 265
the mechanical properties through the TRIP effect. Sev- ation is possible, i.e. as long as the austenite grains are
eral experimental [7,34] and now theoretical [35] studies larger than the platelets length. By contrast, the bainite
showed that the austenite carbon content must lie morphology is completely different when the austenite
within the given range for the right effectiveness of the grain size is only a few micrometer as in the case of
mechanically-activated martensitic transformation. As the intercritical austenite shown in Fig. 4(b). The bainite
a consequence, carbon enrichment also delineates the that forms in very small austenite grains presents adja-
temperature range in which the bainite transformation cent platelets that completely cross the austenite grain.
can be conducted as a function of the T0 curve and thus These differences in the bainite morphology as a func-
of the chemical composition of the grade. tion of the austenite grain size influence in a large way
Beside specific chemical compositions, another fea- the bainite transformation kinetics. Fig. 5 presents the
ture highlight the bainite transformation occurring in evolution of the normalised bainite content as a function
the TRIP-aided steels, the austenite grain size. Indeed, of the transformation time for several austenite grain
as shown in Fig. 1, the austenite grains classically pres- sizes. For smaller grains, the transformation starts ear-
ent a very fine size of the order of 1 or 2 lm. Fig. 4 com- lier but proceeds at a slower rate. Indeed, the reduction
pares the bainite morphology as a function of the size of of the grain size brings about an increase of the grain
the austenite grain in which it occurs. The classical bai- boundary area that accelerates the rate of transforma-
nite sheaf structure can be clearly seen on Fig. 4(a). The tion thanks to an enhanced nucleation rate. This influ-
first bainitic ferrite sub-units nucleate at the austenite ence was already shown and modelled by Rees and
grain boundary and grow towards the interior of the Bhadeshia [36]. However, when the austenite grain size
austenite grain. New sub-units then nucleate and grow is reduced to the length of one platelet, nucleation and
from the tip of the previous ones, bringing about the growth of the next platelet at the tip of the previous
sheaf structure. This process is valid as long as tip nucle- one are no more possible. As a consequence, even if
Fig. 4. TEM micrographs and schematic representation of the growth process and resulting microstructures of bainite in the case of large (a) and
small (b) austenite grains.
P.J. Jacques / Current Opinion in Solid State and Materials Science 8 (2004) 259 265 263
� - � TRIP1 VłR TRIP1
Vąb/Vł0
True Stress
Austenite Volume
� - � TRIP2 VłR TRIP2
(MPa)
Fraction (%)
1.0
� - � TRIP3 VłR TRIP3
1200
0.8
1000
800
0.6 20
600
15
400
10
2 m
�
0.4
200
�
8 m 5
�
50 m
0
0
0.2
0.00 0.05 0.10 0.15 0.20 0.25 0.30
True Strain
Fig. 6. True stress true strain curves and evolution of the retained
0.0
austenite content during plastic straining of several TRIP-assisted
1 10 100 1000
multiphase steels.
Bainitic Holding Time (s)
Fig. 5. Evolution of the normalised bainite volume fraction (bainite
content with respect to the initial amount of austenite) as a function of
typical true stress true strain tensile curves of TRIP-
the isothermal holding time for several austenite grain sizes.
aided steels together with the evolution of the retained
austenite content with true strain. These specimens pres-
the transformation starts earlier, the rate of transforma- ent different initial amounts of retained austenite with
tion is slower. The kinetics of bainite transformation in distinct C enrichments. This Figure clearly shows that
very small grains are therefore controlled only by the the best strength ductility balance occurs for the speci-
grain boundary nucleation rate. men presenting the largest TRIP effect distributed uni-
formly all along plastic straining.
Furthermore, the occurrence of the TRIP effect for
3. Mechanical properties of TRIP-assisted multiphase small austenite grains dispersed in a soft ferritic matrix
steels has a large influence on the work hardening rate. Due
to the shape and volume changes accompanying the
As already mentioned, the objective at the origin of transformation of austenite to martensite, local plastic-
the development of the TRIP-assisted multiphase steels ity is generated in the surrounding ferrite grains [34].
is the improvement of the strength level without sacrific- The TEM micrograph of Fig. 7 illustrates the numerous
ing the uniform elongation. However, these two proper- accommodation dislocations generated within ferrite at
ties are quite antagonist ones. It is indeed very difficult the tip of the deformation-induced martensitic variants.
to increase the stress level that can be sustained by a With respect to the dislocations within the ferrite, the
material without reducing its resistance to the localisa- TRIP effect thus plays the role of an additional source
tion of deformation and resulting fracture. These two that increases the plasticity properties. Quite simple cal-
properties of strength and deformation are actually re- culations [34] showed a clear correspondence between
lated to the work hardening rate, i.e. the increase of the austenite transformation rate and the strength duc-
stress for an increase of strain, dr/de. At the microscopic tility balance of the TRIP-aided steels.
scale, the work hardening rate results from the disloca- However, the increase of the dislocation density can-
tion dynamics, i.e. the balance between the creation not entirely explain the high strength level exhibited by
and annihilation of dislocations. the TRIP-aided steels. Another important aspect of
With respect to other steel grades and microstructures, these steels is the composite nature of their microstruc-
the TRIP-aided steels allow to improve the work harden- tures. They indeed combine phases with antagonistic
ing rate by skillfully combining several mechanisms of properties. Neutron diffraction has allowed the measure-
strengthening and softening. Two mechanisms can effec- ment of the yield strength of the different phases. Values
tively be argued in the case of the TRIP-aided steels: of 500 MPa, 650 MPa, 900 MPa and 2000 MPa were
found for ferrite, bainite, austenite and martensite,
(i) the TRIP effect, respectively [38]. As a result of this large variability of
(ii) the composite-like nature of their microstructures. properties among the phases, stress and strain partition-
ing occurs during loading and dictates the macroscopic
First of all, the TRIP effect has a large influence on the stress strain response [39,40]. Moreover, the present
resulting mechanical properties [34,37]. Fig. 6 presents TRIP-aided steels constitute �evolving composites� since
264 P.J. Jacques / Current Opinion in Solid State and Materials Science 8 (2004) 259 265
Fig. 7. BF and DF TEM micrographs illustrating the dislocations generated in the ferrite at the tip of the strain-induced martensitic variants.
thus corresponds to the strengthening resulting from the
progressive appearance of this martensite.
4. Conclusion
In order to respond to stringent structural applica-
tions, the TRIP-assisted multiphase steels have been
developed through the control of complex phase trans-
formations schemes. Thanks to the diversity of proper-
ties profiles exhibited by the steel phases and the use
of strain-induced phase transformation, unique combi-
nations of strength and deformability can be obtained.
Acknowledgment
The author acknowledges the FNRS and the FRFC
(Belgium). This work was partly supported by the Bel-
gian Science Policy, within the framework of the PAI
Fig. 8. Comparison of the macroscopic stress and the stress calculated
P5/08 project   From microstructure towards plastic
without taking into account the martensitic phase.
behaviour of single- and multiphase materials  .
the proportions of austenite and martensite continu-
References
ously change during straining and definitely modify
the stress partitioning. As a conservative principle, the
[1] Patel JR, Cohen M. Acta Metall 1953;1:531 8.
law of mixture states that the applied stress and strain
[2] Angel T. J Iron Steel Inst 1954:165 74.
are distributed among the different phases with respect
[3] Olson GB, Cohen M. In: Antolovich SC, Ritchie RO, Gerberich
WW, editors. Mechanical properties and phase transformations
to their volume fractions. In the case of stress, it writes
in engineering materials. New Orleans: Met. Soc. AIME; 1986.
0 0
r�e� źfa�a ra�a �e� �fc�e�rc�e� �fa �e�ra �e��1� p. 367 90.
b b
[4] Ludwigson DC, Berger JA. J Iron Steel Inst 1969:63 9.
Thanks to neutron diffraction, it was possible to mea- [5] Magee CL, Paxton HW. Trans Metall Soc AIME 1968;242:
1741 9.
sure the stress level for the BCC (ferrite (a) and bainite
[6] Greenwood GW, Johnson RH. Proc Roy Soc London A
(ab)) and FCC (austenite (c)) phases. Fig. 8 compares
1965;283:403 22.
the measured macroscopic stress level with the calcu-
[7] Jacques PJ. In: Finel A, MaziŁre D, Veron M, editors. Thermody-
lated stress for the 2 first terms of Eq. (1) (i.e. without
namics, microstructures and plasticity. The Netherlands: Kluwer
Acad. Publ.; 2003. p. 241 50.
taking into account of the martensite). The shaded area
P.J. Jacques / Current Opinion in Solid State and Materials Science 8 (2004) 259 265 265
[8] Zaefferer S, Ohlert J, Bleck W. Acta Mater 2004;52:2765 78. [24] Godet S, Harlet P, Delannay F. Mater Sci Forum 2003;426-
[9] Godet S, Georges C, Jacques PJ. In: Damm EB, Merwin MJ, 4:1433 8.
editors. Austenite formation and decomposition. Warrendale, [25] Bhadeshia HKDH, Edmonds DV. Metall Trans A 1979;10A:
Pennsylvania: ISS & TMS; 2003. p. 523 36. 895 905.
[10] Zackay VF, Parker ER, Fahr D, Busch R. Trans Am Soc Met [26] Vanden Eynde X, Servais JP, Lamberigts M. Surf Int Anal
1967;60:252 9. 2003;35:1004 14.
[11] Olson GB, Cohen M. Metall Trans A 1975;6A:791 5. [27] Girault E, Mertens A, Jacques P, Houbaert Y, Verlinden B, Van
[12] Bhandarkar D, Zackay VF, Parker ER. Metall Trans 1972;3: Humbeeck J. Scripta Mater 2001;44:885 92.
2619 31. [28] De Meyer M, Vanderschueren D, De Cooman BC. ISIJ Int
[13] Cai XL, Garatt-Reed AJ, Owen WS. Metall Trans A 1985;16A: 1999;39:813 22.
543 51. [29] Traint S, Pichler A, Hauzenberger K, Stiaszny P, Werner E. Steel
[14] Speich GR, Demarest VA, Miller RL. Metall Trans A 1981;12A: Res Int 2002;73:259 66.
1419 28. [30] Chen HC, Era H, Shimizu M. Metall Trans A 1989;20A:437 45.
[15] Rigsbee JM, VanderArend PJ. In: Davenport AT, editor. Form- [31] Sakuma Y, Matlock DK, Krauss G. Metall Trans A 1992;23A:
able HSLA and dual-phase steels. Warrendale: TMS-AIME 1221 32.
AIME; 1977. p. 56 86. [32] Girault E, Jacques PJ, Ratchev P, Van Humbeeck J, Verlinden B,
[16] Jacques PJ, Girault E, Harlet P, Delannay F. ISIJ Int 2001;41: Aernoudt E. Mater Sci Eng A 1999;273-275:471 4.
1061 7. [33] Jacques PJ. PhD thesis, UCL, 1998.
[17] Jacques PJ, Girault E, Mertens A, Verlinden B, Van Humbeeck J, [34] Jacques PJ, Furn�mont Q, Mertens A, Delannay F. Phil Mag A
Delannay F. ISIJ Int 2001;41:1068 74. 2001;81:1789 812.
[18] Bhadeshia HKDH. Bainite in Steels. 2nd ed. London: IOM; [35] Van Rompaey T, Furnemont Q, Jacques PJ, Pardoen T, Blanpain
2001. pp. 189 200. B, Wollants P. Steel Res Int 2003;74:365 9.
[19] Bhadeshia HKDH, Christian JW. Metall Trans A 1990;21A: [36] Rees GI, Bhadeshia HKDH. Mater Sci Technol 1992;8:985 93.
767 97. [37] Jacques PJ, Bell T. In: Cohen JN, editor. Proc Int Symp Honor
[20] Ohmori Y, Maki T. Mater Trans JIM 1991;32:631 41. of Profs. Materials Park OH: ASM; 2000. p. 539 46.
[21] Takahashi M, Bhadeshia HKDH. Mater Trans JIM 1991;32: [38] Furn�mont Q. PhD thesis, UCL, 2003.
689 96. [39] Furn�mont Q, Lacroix G, Godet S, Conlon KT, Jacques PJ. Can
[22] Jacques PJ, Girault E, Catlin T, Geerlofs N, Kop T, van der Metal Quaterly 2004;43:35 42.
Zwaag S, Delannay F. Mater Sci Eng A 1999;273-275:475 9. [40] Jacques PJ, LadriŁre J, Delannay F. Metall Mater Trans A
[23] Jacques PJ. J Phys 2003;112:297 300. 2001;32A:2759 68.