Conference PaperPDF Available

Development of TBF Steels with 980 MPa Tensile Strength for Automotive Applications: Microstructure and Mechanical Properties

Authors:
Development of TBF Steels with 980 MPa Tensile Strength for Automotive Applications: Microstructure
and Mechanical Properties
A. Bachmaier, K. Hausmann, D. Krizan, A. Pichler
voestalpine Stahl Linz GmbH, voestalpine-Str. 3, 4020 Linz, Austria
Key words: TBF grades, influence of the processing parameters¸ continuous annealing line, mechanical properties, microstructure
ABSTRACT
In this paper, an overview over the development of industrially processed advanced high strength sheet steels for automotive
applications with a tensile strength of 980MPa and an improved ductility compared to the widely used high-strength dual-phase and
complex phase automotive steel grades is given. For these advanced high strength steels with tensile strength above 980MPa which
enable lighter and safer car-bodies a steadily increasing demand is observed. Optimal processing parameters are chosen according to
different heat treatment schedules in lab trials. Phase transformations are investigated by dilatometric measurements. Furthermore, the
mechanical properties are determined. The microstructure of these materials plays a key role in their mechanical properties. Therefore,
the parameters of the heat treatment cycles are adapted to obtain the microstructure which lead to the desired mechanical properties
during processing via continuous annealing lines.
INTRODUCTION
As a response to the steadily increasing requirements of the automotive sector regarding lighter and safer car bodies, advanced high
strength sheet steels with a tensile strength above 980 MPa and improved formability (ductility and bendability) compared to
conventional dual-phase and complex-phase steel grades with the same strength level are currently developed by the steel industry [1].
TRIP bainitic ferrite (TBF) and Quench-and-Partitioning (Q&P) steels are such steel grades which are belonging to the group of so
called “third generation of advanced high strength steels” [2]. The microstructure of TBF steels consists of a matrix of bainitic ferrite
with retained austenite inclusions and is produced by isothermal holding in the bainitic regions after fast cooling from fully austenitic
microstructures [3]. Q&P steels exhibit a matrix which consists of martensite, tempered martensite or/and lower bainite with retained
austenite inclusions produced by isothermal holding after partial martensite transformation by short cooling or by isothermal holding
below the martensite start after fast cooling from fully austenized microstructures [4]. During straining, the TRIP (TRansformation
Induced Plasticity) effect of retained austenite known from conventional TRIP steels is used as well which leads to excellent
formability at room temperature. By the combination of the different phases in TBF and Q&P steels, the properties can be tailored for
specific applications. TBF steels (“dual-phase type” with high strain hardening) which can be used for deep drawing applications or
Q&P steels (“complex-phase type” exhibiting higher yield ratios, reduced elongations and reduced n-values) with an excellent
bendability can be used for bending operations or roll forming.
The most important step during processing is the annealing of the as-cold rolled material via continuous annealing line. Typical “base”
chemical compositions of TBF steels contain C, Si and Mn as major alloying elements. Alloy modifications include variations of the
Al, Nb and Cr content [5]. Si suppresses the formation of cementite during bainitic transformation which enhances the C content in
retained austenite and thus allows the austenite to be stabilized by carbon [6, 7]. High Si contents of 1.5wt% are usually used in these
types of steels [8]. As a consequence, the transformation of retained austenite into martensite upon deformation and/or thermally
produced martensite during final cooling is prevented. Although Si is of major importance to prevent carbide precipitation during
annealing of the cold rolled material, it causes problems during processing via continuous annealing lines. Si alloyed steels exhibit
selective oxidation at the steel surface resulting, for example, in deteriorated galvanizability [9]. Therefore, other alloying elements
having a similar effect of suppressing carbide formation have to be considered. The influence of reduced Si contents (<1wt%) in
combination with the addition of other alloying elements completely or partially substituting Si has been studied recently [10]. Based
on the results, commercially produced cold rolled material was manufactured, investigated in detail by different lab trials to obtain
ideal processing parameters for the final annealing via continuous annealing to obtain TBF and Q&P steel grades and produced on
industrial scale to obtain the desired mechanical properties.
EXPERIMENTAL PROCEDURE
The results in this paper are obtained from a commercially produced cold rolled material which was cold rolled to a final thickness of
1.4 mm. The chemical composition is given in Table 1.
Table 1: Chemical composition of the investigated steel grade in wt. %.
C
Si+Al
Mn+Cr+Mo
Nb
+Ti
TBF steel
0.2
<3
0.05
To optimize annealing parameters and to study the transformation behavior, annealing simulations in the laboratory with the Multi-
Purpose Annealing Simulator (MULTIPAS) and dilatometric investigations on a Bähr dilatometer DIL805 A/D were performed on
specimens prepared from the commercially processed cold rolled material. Dilatometric investigations with different cooling rates and
varying overaging temperatures and time (I-IV) were conducted. Volume fractions of transformed phases during the dilatometric
investigations are derived from the dilatometric changes using the lever rule.
Fig.1: Schematics of the applied laboratory annealing series.
Furthermore, annealing series on the basis of actual continuous annealing line layouts were conducted in the laboratory with the
MULTIPAS (III-IV). Summarized, the following annealing series were applied:
I) Specimens were cooled with different cooling rates (3K/s - 80K/s) to room temperature after annealing for 60s at 900°C.
II) Specimens were heated to an annealing temperature of 900°C, held at the annealing temperature for 60s to obtain fully austenitic
microstructures and subsequently quenched with a cooling rate of 70K/s to different isothermal holding temperatures between 325°C
and 500°C for an isothermal holding time of 600s.
III) Specimens were heated to an annealing temperature of 850°C, held at the annealing temperature for 60s to obtain fully austenitic
microstructures and subsequently quenched with a cooling rate of 30K/s and 50K/s to three isothermal holding temperatures (400°C,
425°C, and 450°C) and for different isothermal holding times of 30-600s. Dilatometric investigations were only performed with an
overaging time of 600s.
IV) Q&P heat treatment cycles according to industrial processing and annealing layouts were performed. Specimens were heated to
annealing temperatures of 850°C, held for 60s, cooled with two different cooling rates (30K/s and 50K/s) to three different quench
temperatures (350°C- 380°C) and two different overaging temperatures (400°C and 440°C).
A schematic of the annealing series applied in dilatometric investigations as well as in the annealing simulations is shown in Fig. 1.
The most important annealing parameters for the applied annealing series are summarized in Table 2.
Samples for light optical metallography were prepared by standard metallographic preparation (etched with LePera’s etchant) [11].
Volume fraction of retained austenite was measured using the saturation magnetization method [12]. The mechanical properties with
tensile specimens machined with their tensile axis parallel to the rolling direction were determined according to testing procedure DIN
EN ISO 6892-1.
Additionally, cold rolled sheets were industrially processed via continuous annealing line by choosing the optimal processing
parameters obtained by the lab trials and subsequent microstructural characterization was performed as described above. Furthermore,
the mechanical properties with tensile specimens machined with their tensile axis parallel and perpendicular to the rolling direction
were determined.
Table 2: Annealing parameters (Tan...annealing temperature, tan…annealing time, TOA…overaging temperature, tOA…overaging time,
TQ…quenching temperature and CR…cooling rate) for dilatometric investigations (I-IV) and MULTIPAS simulations (III-IV).
Annealing series T
an
(°C) t
an
(s) CR (K/s) T
Q
(°C) T
OA
(°C) t
OA
(s)
I
900
60
3
-
80
-
-
-
II
900
60
7
0
-
325
-
500
600
III 850 60
30
-
400/425/450 30/120/300/600
50
-
IV 850 60
30
350/360/380
400/440
600
50
350/360/380
400/440
600
RESULTS AND DISCUSSION
Lab trials
In order to optimize annealing conditions to obtain TBF and Q&P steel grades, dilatometric measurements were performed. The onset
of austenite formation during heating (Ac1=753°C), the temperature of the completion of austenite formation (Ac3=831°C) and the
martensite transformation start and finish temperature (Ms=371°C and Mf=289°C, respectively) were calculated from dilatometric data
during continuous cooling (I). Furthermore, the isothermal transformation behavior in the bainitic holding range was investigated
which is an essential step during the production of TBF grades via continuous annealing lines. The overaging temperature influences
not only the transformation behavior and the kinetics of transformation, it also determined the amount of austenite transformed during
isothermal holding. A phase transformation during cooling from the annealing temperature could not be observed for cooling rates
>50 K/s in the continuous cooling experiments. Therefore, a cooling rate of 70K/s and holding temperatures well above Ms were
chosen to solely investigate the transformation behavior in the bainitic formation range (II). Fig.2a shows the transformation kinetics
of the austenite during the isothermal holding at different holding temperatures from 400°C to 500°C. For the highest isothermal
holding temperature a two-step reaction is observed. The first step of transformation which is the bainitic reaction is completed very
fast (tOA<60s). Carbide precipitation is subsequently observed. At lower overaging temperatures, a one-step transformation behavior is
observed but the transformation kinetic is decelerated with decreasing overaging temperature due to a lower diffusion. Nevertheless an
increasing amount of bainite is formed with decreasing overaging temperature. Transformation kinetics is rather slow at 400°C and
425°C, but the transformation is completed within the overaging times which are typical of a continuous annealing line (~600s) and
the overall amount of formed bainite is maximized. The temperature dependence of the bainite formation can be explained by the
well-known To-concept [13]. If a displacive growth mechanism of bainitic ferrite is assumed, the transformation stops if the C content
of the remaining austenite reaches the To-boundary which is given by the intersection point between Gibb’s energy curves of ferrite
and austenite having an identical composition at a certain temperature. Lower overaging temperatures allow higher C contents in the
untransformed austenite. For transformation temperatures of 400°C and 425°C, carbide precipitation is not observed within the
holding times investigated and high amounts of retained austenite are stable at room temperature. At higher holding temperatures
(450°C and 500°C) the amount of austenite which transforms into bainite becomes less. If carbide precipitation additionally occurs
during isothermal holding (transformation temperature of 500°C), the austenite transforms nearly completely to martensite during final
cooling to room temperature and no or very less retained austenite is stabilized at room temperature which is shown in Fig.2b.
Fig.2: a) Dilatation-time curves obtained from isothermal holding and b) content of retained austenite as a function of the different
overaging temperatures.
In subsequent annealing simulations on the basis of actual continuous annealing line layouts (III), the influence of the overaging
temperature (400-450°C) and time in the bainitic range on the microstructure and mechanical properties with two different cooling
rates (30K/s and 50K/s) was studied (Fig.3a-c). Comparing the different cooling rates, yield and tensile strengths are slightly enhanced
whereas the uniform elongation values are lower for higher cooling rates (Fig.3a-b). Faster cooling rates lead to lower amounts of pro-
eutectoid ferrite and/or bainitic ferrite during cooling. In Fig.3d, the transformed phase fractions before reaching the overaging
temperature for both cooling rates as a function of temperature are plotted. A significant higher amount of transformation at a cooling
rate of 30K/s before reaching the final overaging temperature results in lower measured tensile strength values. In general, a cooling
rate of 30 K/s provides a better combination of strength and elongation values. Comparing different overaging temperatures at this
cooling rate, there is no significant influence of the overaging temperature on the tensile strength for low bainitic temperatures (400°C
or 425°C) if the overaging time is >300s. For short overaging times (<300s), an influence of the slower transformation kinetics at
lower overaging temperatures shown in Fig.2a are reflected in the mechanical properties as well. The tensile strength at these
overaging temperatures decreases due to the increasing amount of bainite formed during longer isothermal holding at the overaging
temperature. As a consequence, the stability of retained austenite is increased which leads to a lower amount of retained austenite
which transforms into martensite upon deformation and/or thermally produced martensite during final cooling. Due to the fastest
transformation kinetics and the lowest amount of bainite formed at an overaging temperature of 450°C, a certain amount of retained
austenite transforms to martensite during final cooling and the highest tensile strength values are measured of specimens annealed at
this overaging temperature which are not influenced by the overaging time.
Yield strengths are significantly influenced by the overaging temperatures whereby increasing amounts of bainite formed results in
increasing yield strength values. Only the specimens annealed at 400°C exhibit higher yield strengths than 700 MPa. The lowest yield
strength are obtained for an overaging temperature of 450°C independent of overaging time which fits quite well to the transformation
behavior observed.
A maximum of the uniform elongation is observed for each overaging temperature. For overaging temperatures of 425°C and 450°C,
the highest uniform elongation values are obtained at an overaging time of 120s. At the lowest overaging temperature, the maximum is
shifted to longer holding times. At high overaging times, nearly no difference between uniform elongation values for the samples
annealed at 450°C and 425°C with a cooling rate of 30K/s is visible. A higher amount of retained austenite stabilized at room
temperature results in higher uniform elongations values. The stability of retained austenite is a combined effect of a chemical
stabilization due to an optimum concentration of carbon in austenite and a size effect [14, 15]. At high overaging temperatures,
austenite to bainite transformation is quite fast but the austenite carbon concentration is reduced for longer overaging times by carbide
precipitation which lowers the stability of the retained austenite. Due to the slower kinetics of austenite transformation at lower
overaging temperature, application of longer overaging times results in higher uniform elongation values. An ideal combination of
yield strength, tensile strength and uniform elongation for a TBF steel is obtained for a cooling rate of 30K/s and an overaging
temperature of 400°C.
Fig.3: a-c) Influence of the overaging temperature and overaging time as well as the cooling rate on the mechanical properties. d)
Influence on the cooling rate on the transformed phase fraction upon cooling from the annealing temperature until the final overaging
temperature is reached.
Additionally, Q&P heat treatment annealing simulations according to industrial processing and annealing layouts were performed to
obtain parameters for industrial processing of Q&P steels (IV). The influence of three different quench temperatures (350°C, 360°C
and 380°C) on the microstructure and mechanical properties with two different cooling rates (30K/s and 50K/s) and two different
overaging temperatures was studied (Fig.4a-d). Two quench temperatures were chosen to be under and one quench temperature to be
slightly above the martensite transformation temperature (Ms=371°C) determined from continuous cooling experiments. In general,
higher yield and tensile strength values are obtained at the higher cooling rate which are steadily increasing with decreasing overaging
temperature. The cooling rate determines the tensile strength values which are nearly unaffected from the different overaging
temperatures. In contrast, the yield strength values are significantly reduced at higher overaging temperatures. The overaging
temperature has also a certain influence on the uniform and total elongation which is enhanced for higher overaging temperatures.
Regarding the influence of the cooling rate, higher values for uniform and total elongation are obtained for the lower cooling rate.
In Fig.5, transformation maps illustrating the transformed phase fractions during annealing obtained from dilatometric investigations
with the same annealing parameters used for the annealing simulations described above are shown. The map provides an overview
about the phase transformations which occur during the heat treatment cycle. Before the final quench temperature is reached the
occurring phase transformations are related to ferrite or bainitic ferrite formed during cooling. Phase transformations directly at the
quench temperature correspond to the athermal formation of martensite which starts rapidly if the quenching temperatures are below
the martensite start temperature. Phase transformations during isothermal holding at the overaging temperature are referred to the
formation of bainite. Very small additional dilation of the dilatometric specimen can be further caused by C partitioning from
martensite/bainite to austenite during isothermal holding at the overaging temperature [16]. Furthermore, the amount of measured
retained austenite at room temperature is given in the maps.
Depending on the final quench temperature, up to 8% more phase transformation occurs until the quench temperature is reached at the
slower cooling rate. As a consequence, the amount of transformed phase fractions directly at the quench temperature is significantly
decreased. The cooling rate has also a small influence on the amount of retained austenite which is slightly higher at lower cooling
rates.
Fig.4: Influence of the quenching temperature and overaging temperature as well as the cooling rate on the mechanical properties.
Yield and tensile strength for cooling rates of 30K/s and 50K/s for a) an overaging temperature of 400°C and b) 440°C. Uniform and
total elongation for cooling rates of 30K/s and 50K/s for c) an overaging temperature of 400°C and d) 440°C.
Assuming that the transformed phase during cooling is ferrite and/or bainitic ferrite and martensite at the quench temperature, the
lower yield and tensile strength values as well as the higher uniform and total elongation values at the slower cooling rate can be
explained by the lower fraction of “hard” martensitic phase and higher fraction of “soft” ferrite and/or bainitic ferrite phase formed
during cooling in the microstructure. The significantly lower yield strengths obtained at higher overaging temperatures for both
cooling rates are a consequence of the higher fraction of upper bainite which is obtained by transformation at high overaging
temperatures. Higher uniform and total elongation values at the higher overaging temperature are due to enhanced amounts of retained
austenite and due increased amount of upper bainite in the microstructure as well.
Comparing the tensile strength differences between both overaging temperatures at a certain cooling rate, the difference is nearly
constant for all quenching temperatures and the strength can be mainly controlled by choosing a suitable amount of ferrite and/or
bainitic ferrite and martensite. Overaging temperature differences of 40°C between an overaging temperature of 400°C and 440°C are
obviously too small to result in a significant softening effect due to tempered martensite. Although higher elongation values are
obtained for higher overaging temperatures, yield strength values are below 800 MPa. Good combination of yield strength, tensile
strength and uniform elongation for a Q&P steel is given at a cooling rate of 30K/s, the entire range of quench temperatures
investigated and an overaging temperature of 400°C.
Fig.5: Transformation maps obtained from dilatometric investigations for an overaging temperature of 400°C for a cooling rate of
30K/s (a) and 50K/s (b) as well as for an overaging temperature of 440°C for a cooling rate of 30K/s (c) and 50K/s (d). Phase
transformations before reaching the quench temperature (Ferrite and/or bainite formed during cooling), at the quench temperature
(Martensite formed during cooling) and during isothermal holding (Bainite formed during austempering) are shown. The amount of
retained austenite (ret) is also given.
Production on industrial scale
Based on the results from the annealing simulations and dilatometric measurements on industrially cold-rolled material the optimal
annealing parameters were selected to produce TBF and Q&P steel grades. An overaging temperature of 400°C provides an optimal
combination of sufficiently fast transformation kinetics with an optimal amount of bainite formed to obtain the desired mechanical
properties for TBF steels (yield strength>700 MPa, tensile strength>980 MPa, total elongation>14%). The microstructure of the
continuously annealed TBF steel is shown in Fig.6a which consists of bainite and retained austenite. The amount of retained austenite
is ~12%. Typical mechanical properties of the continuously annealed TBF steel which are obtained from longitudinal (L) and
transversal (T) specimens are shown in Fig.6b. The yield strength and the tensile strength are about 762-771 MPa and 1021-1024
MPa, respectively. The resulting yield strength ratio is, therefore, ~0.75. The values of the total elongation are about 15%,
respectively. At strain levels between 4 and 6%, the n-value is 0.16. n-values about 0.14 are measured at strain levels between 6 and
10 %. The excellent properties can be also described by the products of tensile strength times total elongation which is about 15.300
MPa% for the TBF steel grade.
Fig. 6: a) Microstructure and b) obtained mechanical properties of a continuously annealed TBF steel obtained from longitudinal (L)
and transversal (T) specimens.
A quench temperature below 380°C and an overaging temperature of 400°C provides an optimal combination of an optimal amount of
martensite and of bainite formed during the isothermal overaging to obtain the desired mechanical properties for Q&P steels (yield
strength>800 MPa, tensile strength>980 MPa, total elongation>10%). The microstructure of the continuously annealed Q&P steel is
shown in Fig.7a which consists of lower bainite, small fractions of martensite and retained austenite. The amount of retained austenite
is ~9%. Typical mechanical properties obtained from longitudinal (L) and transversal (T) specimens are shown in Fig.7b.
Fig. 7: a) Microstructure and b) obtained mechanical properties of a continuously annealed Q&P steel obtained from longitudinal (L)
and transversal (T) specimens.
The yield strength and the tensile strength are about 870-880 MPa and 1080-1090 MPa, respectively. The resulting yield strength ratio
is, therefore, ~0.81. The values of the total elongation are about 11-12 %. At strain levels between 2 and 4%, the n-value is 0.12. n-
values about 0.13 are measured at strain levels between 4 and 6 %. The excellent properties can be also described by the products of
tensile strength times total elongation which is about 13.000 MPa%.
SUMMARY
The influence of typical continuous annealing line processing parameters on the microstructure and mechanical properties of advanced
high strength steel grades was investigated for industrially processed cold-rolled material with a Si content <1wt%. Based on
experimental results optimized annealing cycles were selected to successfully industrially produce TBF steel grades with a minimum
yield strength of 700 MPa (minimum tensile strength 980 MPa) as well as Q&P steel grades with a minimum yield strength of 800
MPa (minimum tensile strength 980 MPa) via continuous annealing line. The total elongations are clearly above 14% and 10% for
TBF steel grades and Q&P steel grades, respectively. Compared to conventional dual-phase and complex phase automotive steel
grades with tensile strengths of 980 MPa, an enhanced ductility is reached. The strength and ductility can be simply varied by
adjusting the amount of “hard” phase i.e. martensite or lower bainite in combination with suitable amounts of “soft” phase i.e. ferrite,
upper bainite and retained austenite due to appropriate annealing conditions.
REFERENCES
1. L. Samek and D. Krizan, “Steel – Material of choice for automotive lightweight applications”, Conference Proceeding of
Metal 2013, Brno, Czech Republic, 2013.
2. E. De Moor, P.J. Gibbs, J.G. Speer and D.K. Matlock, “Strategies for third-generation advanced high-strength steel
development”, Iron and Steel Technology, 7/11, 2010, pp.133-144.
3. K. Sugimoto, M. Tsznezawa, T. Hojo and T. Ikeda, “Ductility of 0.1-0.6C-1.5Si-1.5Mn ultra high strength TRIP aided sheet
steel with bainitic ferrite matrix”, ISIJ International, Vol.44, 9, 2004, pp.1608-1614.
4. B.C. De Cooman and J. Speer, “Quench and partitioning steel: New AHSS concept for automotive anti-intrusion
applications”, Steel Research International, Vol. 77, 9-10, 2006, pp.364-640.
5. S. Traint, A. Pichler, R. Tikal, P. Stiaszny and E.A. Werner, Influence of manganese, silicon and aluminium on the
transformation behavior of low-alloyed TRIP steels”, 42nd Mechanical Working and Steel Processing Conference, Toronto,
2000.
6. E. Kozeschnik and H. K. D. H. Bhadeshia, “Influence of Silicon on Cementite Precipitation in Steels”, Materials Science and
Technology, Vol. 24, 2008, pp.343-347.
7. H. K. D. H. Bhadeshia, “The Bainite Transformation: Unresolved Issues”, Materials Science and Engineering A, Vol. A273-
275, 1999, pp.58-66.
8. S.A. Khan and H.K.D.H. Bhadeshia, “Kinetics of Martensitic transformation in partially bainitic 300M steel”, Materials
Science and Engineering A 129, 1990, pp.257-272.
9. J. Mahieu, S. Claessens and B.C. Cooman, “Galvanizability of High-Strength steels for automotive applications”,
Metallurgical and Materials Transactions A, Vol.32A, 2001, pp.2905-2908.
10. K. Hausmann, D. Krizan, K. Spiradek-Hahn, A. Pichler and E. Werner, Steel Research International, in preparation.
11. F.S. LePera, Journal of Metals, 1980, pp. 32-38.
12. E. Wirth, A. Pichler, R. Angerer, P. Stiaszny, K. Hauzenberger, Y.F. Titovets and M. Hackl, “Determination of the volume
amount of retained austenite and ferrite in small specimens by magnetic measurements”, International Conference on TRIP-
Aided High Strength Ferrous Alloys, Ghent, Belgium, June 2002, pp.61-64.
13. H. K. D. H. Bhadeshia, “Bainite: The Incomplete Reaction Phenomenon and the Approach to Equilibrium, Proceedings
Conference Solid Phase Transformations”, Pittsburgh, USA, 1981, pp.1041-1048.
14. P J Jacques, “Transformation-induced plasticity for high strength formable steels”, Current Opinion in Solid State &
Materials Science, 8(3-4), 2004, pp.259–265.
15. J.Wang and S. Van Der Zwaag, “Stabilization mechanisms of retained austenite intransformation-induced plasticity steel”,
Metallurgical and Materials Transactions A-Physical Metallurgy and Materials Science, Vol.32A, 2001, pp.1527–1539.
16. M.J. Santofimia, L. Zhao, R. Petrov, C. Kwakenraak, W.G. Sloof and J. Sietsma, “Microstructural development during the
quenching and partitioning process in newly designed low-carbon steel”, Acta Materialia, Vol.59, 2011, pp.6059-6068.
... cooling from an austenite microstructure into the bainitic regions followed by isothermal holding. According to [21] these steels are also named as DH steels, dual phase steels with higher global ductility/formability. The microstructure of Q&P steels consists of martensite, tempered martensite and bainite with retained austenite. ...
... The underlying phase-illustrations of the microstructure are presented in Table A2. The grades with higher global formability display a high amount of retained austenite, which is responsible for the "TRIP"-effect during loading [21]. The amounts lay in a range from 2.6% (CH1200) up to 10.7% (DH1200). ...
... Noticeable in the SEM micrographs of both materials, displayed in Table A19, is the occurrence of bigger voids for the DH grade, which is similar for all DH-and CH-grades and might result from a damage mechanism related to the retained austenite. The global formability of the DH800 is on nearly the same level as the DP600 s, because of the "TRIP"-effect [21] during loading related to the retained austenite. The local formability remains on nearly the same level compared to the DP800's. ...
Article
Full-text available
The usage of high-strength steels for structural components and reinforcement parts is inevitable for modern car-body manufacture in reaching lightweight design as well as increasing passive safety. Depending on their microstructure these steels show differing damage mechanisms and various mechanical properties which cannot be classified comprehensively via classical uniaxial tensile testing. In this research, damage initiation, evolution and final material failure are characterized for commercially produced complex-phase (CP) and dual-phase (DP) steels in a strength range between 600 and 1000 MPa. Based on these investigations CP steels with their homogeneous microstructure are characterized as damage tolerant and hence less edge-crack sensitive than DP steels. As final fracture occurs after a combination of ductile damage evolution and local shear band localization in ferrite grains at a characteristic thickness strain, this strain measure is introduced as a new parameter for local formability. In terms of global formability DP steels display advantages because of their microstructural composition of soft ferrite matrix including hard martensite particles. Combining true uniform elongation as a measure for global formability with the true thickness strain at fracture for local formability the mechanical material response can be assessed on basis of uniaxial tensile testing incorporating all microstructural characteristics on a macroscopic scale. Based on these findings a new classification scheme for the recently developed high-strength multiphase steels with significantly better formability resulting of complex underlying microstructures is introduced. The scheme overcomes the steel designations using microstructural concepts, which provide no information about design and production properties.
... A TBF-acélok tipikus kémiai összetétele C, Si és Mn fő ötvözőelemeket tartalmaz. További szokásos ötvözők az Al, Nb és Cr különböző összetétel-kombinációkban [13]. A Si gátolja a karbid képződését a bainites fázisátalakulás során, ami növeli a maradék ausztenit C-tartalmát, és ezáltal lehetővé teszi a maradék ausztenit karbonnal való stabilizálását. ...
Article
Az utóbbi évtizedekben az autóipar folyamatosan növekvő követelményekkel nézett szembe. A növekvő követelmények mind a felhasználók, mind pedig a jogi előírások terén megfigyelhetők. Fogyasztói oldalról az egyre gazdaságosabb, kisebb fogyasztású, ugyanakkor magasabb kényelmi szintet és nagyobb biztonságot adó személygépkocsik iránti igény jelentkezik. Ezeket a követelményeket kiegészítik a fokozott környezetvédelmi előírások, amelyek minél kisebb károsanyag-kibocsátást tesznek lehetővé. Ezeknek a gyakran egymásnak is ellentmondó követelményeknek a kielégítése hatalmas kihívást jelent az autógyártókkal és az alapanyagot gyártókkal szemben. A követelmények kielégítésének igénye az elmúlt 40–50 évben hatalmas fejlődést eredményezett az autógyártóknál és az alapanyaggyártóknál is. Ebben a cikkben röviden áttekintjük az autóiparral szemben támasztott fő követelményeket, amelyek egyben az alapanyag-fejlesztés legfontosabb hajtóerőit is jelentik. A hagyományos nagy szilárdságú acélok mellett elsősorban a korszerű nagy szilárdságú acélok három generációjának ismertetésével foglalkozunk.
... Tensile strength levels for Q & P and TBF steels are usually between 1,000 and 1,500 MPa, with total elongation between 10% and 20%. Bachmaier et al. [10] have investigated the influence of typical continuous annealing line processing parameters on the microstructure and mechanical properties of the industrially produced TBF and Q&P steels with a minimum tensile strength of 980 MPa. Figure 4. shows the spider diagrams of the mechanical properties obtained from the experimental results using a tensile test for specimens with a gauge length of 80 mm measured in longitudinal and transversal direction. It can be claimed that Q&P steel characterized by a lower total elongation and higher yield strength compared to the TBF steel, which means both steels could be used in various structural and safety automobile co- ...
Article
Full-text available
The modern automobile industry reached a high level of development of steel grades with extraordinary strength and formability, through different new production processes and new processing strategies. The third generation of advanced high strength steel is a new concept that have been developed lately and have already been incorporated in the modern car body structures. The main objective of this work is to review some recent third-generation advanced high strength steel grades applied in sheet metal forming processes mainly from the viewpoint of automotive industry.
... These steels are essentially bainitic ferrite matrix with metastable retained austenite inclusions. TBF chemical composition is containing (0.15 wt% C), (1.5 wt% Si), (1.5 wt% Mn), in addition to alloy modifications comprises Al, Nb and Cr content [5]. TBF steels produced by hot and cold rolling processes [6]. ...
Article
Full-text available
The modern vehicles demand a better fuel economy, decrease in ozone harming substance outflows, and superior safety requirements led to new developments of steel grades with higher strength and good formability. Third generation of advanced high strength steels are the next stage for the automotive companies in steel sheets development. The principal concept of third generation of AHSS is to reap the mechanical properties benefits from first and second generation of AHSS at cost neither too high nor too low. This literature review summarizes the results achieved in a previous paper of the Third Generation of Advanced High Strength Sheet steels literature published by D. Krizan et al. Where we intend to focus on, the recent developments and future trends of the third generation of advanced high strength sheet steels (3-GEN AHSSs) including quenching and partitioning (Q&P), TRIP bainitic ferrite (TBF), medium manganese, density reduced TRIP (δ-TRIP) and nano steels for the modern automotive industry, with emphasis on their main characteristics, processing, and applications .
... Up to now, TBF chemical composition is containing (0.15 wt % C), (1.5 wt % Si), (1.5 wt % Mn), in addition to alloy modifications comprises Al, Nb and Cr content (Bachmaier et al., 2013). TBF steels are produced by hot and cold rolling processes. ...
Conference Paper
Full-text available
The modern automobile industry reached a high level of development of steel grades with extraordinary strength and formability, as a result of the increasing demand for minimizing the fuel economy, decreasing in harming substance outflows, and improve safety precautions. The third generation of advanced high strength steel is a further development of the previous first and second generations for the vehicle companies in steel sheets. Furthermore, the transformation-induced plasticity (TRIP) effect playing an important role in this generation to achieve an excellent combination of strength and ductility. The third generation of AHSS was created and developed in order to have a new grade of steels with superior mechanical properties at a reasonable cost. The current paper reviews the main types of this new generations including quenching and partitioning (Q&P), TRIP bainitic ferrite (TBF), medium manganese (Medium-Mn), density reduced TRIP (δ-TRIP) and nano steels for the modern automotive industry, besides the TRIP effect will be illustrated in detail.
... Typical chemical compositions of TBF steels contain C, Si and Mn as major alloying elements. Alloy modifications include variations of the Al, Nb and Cr content [23]. The cementite formation during bainitic transformation is suppressed by the Si constituent. ...
Chapter
Full-text available
Sheet metal forming is one in all the foremost important production processes in car manufacturing; therefore its developments are significantly determined by the demands of the automotive industry. Recent trends in car production are also characterized by applying lightweight principles. Its main priority is to fulfil both the customers’ demands and also the increased legal requirements. Applying high strength steels could also be thought to be one in all the potential possibilities. Applying high strength steels have a positive response for several of the requirements: increasing the strength may result in the appliance of thinner sheets leading to significant mass reduction. Mass reduction ends up in lower consumption and increased environmental protection. Increasing strength often leads to a decrease in formability. In this paper, an outline of recent material developments within the automotive industry concerning the employment of recent generation advanced high strength steels are going to be given.
... Typical chemical compositions of TBF steels contain C, Si, and Mn as major alloying elements. Alloy modifications include variations of the Al, Nb, and Cr content [47]. The cementite formation during bainitic transformation is suppressed by the Si constituent. ...
Chapter
Full-text available
The automotive industry plays a determinant role in the economy of developed countries. Sheet metal forming is one of the most important processes in car manufacturing. Recent trends in car production may be characterized by the application of lightweight principles. Its main priority is to fulfill both the customers' demands and the increased legal requirements. The application of high strength steels may be regarded as one of the potential possibilities. Applying high strength steels has a positive response for many of the requirements: increasing the strength may lead to the application of thinner sheets resulting in significant mass reduction. Mass reduction is leading to lower consumption with increased environment protection. However, increasing the strength can often lead to the decrease of formability, which is very unfavorable for the forming processes. In this chapter, an overview of recent material developments in the automotive industry concerning the use of new-generation high strength steels will be given. In this paper, the material developments are emphasized from the point of view sheet metal forming; therefore, our focus is on the body-in-white manufacturing in the automotive industry.
Article
A new steel chemical composition is combined with a new press hardening process, in which die-quenching is interrupted by opening the forming tool to permit slow cooling of the hot formed part through the anisothermal bainitic ferrite transformation. This promotes carbon partitioning to austenite before the forming tool is re-closed and die-quenching is resumed to near-ambient temperature. The final microstructure is predominantly bainitic ferrite with dispersions of martensite and up to 11 % retained austenite. Retained austenite can undergo stress induced transformation to martensite in an automobile crash event. The steel exhibits up to 25 % elongation and 930 MPa tensile strength. In contrast to traditional cold formable Transformation Induced Plasticity assisted steels, where retained austenite is consumed during work hardening of cold forming, here, the desired microstructure is achieved after hot forming meaning the retained austenite is more uniformly distributed within the formed part, which enhances energy absorption. The new steel chemical composition is carefully designed to provide optimal microstructural evolution within the constraints of the new press hardening process, yet relatively lean and manufacturer friendly. The new press hardening process is energy efficient as secondary heating is not required since retarded cooling through the bainitic ferrite transformation is provided by residual heat accumulation of the newly developed titanium alloy forming tool. Development of the new technology is demonstrated by press hardening experiments, tensile testing, microstructural analysis, transversal & axial crush testing of formed parts and numerical simulation of crush testing, including a new modelling technique that more accurately simulates deformation of hot versus cold formed parts. Results show a 22 % increase to energy absorption under axial crushing compared to traditional cold formed Transformation Induced Plasticity assisted steels owing to greater work hardening capacity in formed radii of the part, which are shown to be exposed to the highest stresses during crushing.
Conference Paper
Full-text available
A weight reduction, driven by constantly more stringent CO2 regulations along with an improved crashworthiness of modern vehicles, motivates the steel industry to develop novel advanced high strength steels (AHSS), which combine simultaneous increase of both rather contradictory properties: strength and ductility. The steel industry is currently associated with the development of the third generation (3-GEN) AHSS, where the excellent combination of strength and ductility can be achieved by an employment of the transformation-induced plasticity (TRIP) effect - the transformation of a large amount of metastable retained austenite to martensite during straining. This beneficial phenomenon takes place during forming operations with a global ductility such as deep drawing. Apart from that, an excellent performance of the steel during forming operations with a local ductility (e.g. sheet cutting, hole-expansion etc.) becomes also pre-requisite when forming complex components for the automotive industry. Furthermore, a component weight saving via steel density reduction may also come into play, while considering the development of these steel grades. The present paper reviews the above-mentioned aspects resulting in the development of TRIP bainitic ferrite (TBF), quenching and partitioning (Q&P), medium manganese and density reduced TRIP steels for the application in modern automotive platforms. These new developments are thoroughly described from their processing, microstructure and resulting mechanical properties point of view.
Conference Paper
Full-text available
In recent years, for reasons of improved passenger comfort and safety, the weight of passenger cars has continuously increased, leading to higher fuel consumption and greenhouse gas emissions. Since the early 90s competition for safer, lighter, and more fuel economic ground transportation vehicles, triggered by stringent OEMs and governmental demands, led over the years to the market entry of new materials such as new high-strength steels, polymer composites, aluminium alloys, or also magnesium. In this challenging contest of achieving significant emission reductions and fuel economy across all new generation vehicle platforms, the steel industry is today accelerating the implementation of new innovative steels over other emerging materials. The aim of the present contribution is to review the development of novel high strength steels for automotive applications and to highlight their benefits compared to aluminium alloys, as one of their competitor for the automotive lightweight design.
Article
Full-text available
The effects of heat treatment and forming conditions on retained austenite characteristics and ductility of 0.1-0.6C-1.5Si-1.5Mn, mass%, ultra high-strength TRIP-aided sheet steels with bainitic ferrite matrix were investigated. These steels possessed large total elongations of about 20-25% in a tensile strength ranging from 700 to 1 300 MPa when austempered at temperatures above martensite-start temperature (M S). The total elongations were enhanced by warm forming at two temperatures, T P1 and T P2. The first peak forming temperatures T P1s were between 0°C and 75°C and were nearly constant regardless of carbon content of the steels. This was associated with the strain-induced martensite transformation of a large amount of metastable retained austenite which suppressed a rapid fall of strain-hardening rate in an early strain range to resultantly increase the uniform and total elongations. On the other hand, the second peak forming temperatures T P2s were between 200 and 300°C and further large total elongations beyond 30% were achieved in high carbon steels (0.4% C and 0.6% C steels) with tensile strength of 1 300-1 500MPa. The large improvement was controlled by both the strain-induced bainite transformation and dynamic strain aging.
Article
Full-text available
The mechanism by which relatively small concentrations of silicon influence the precipitation of cementite from carbon supersaturated austenite and ferrite are investigated. It is found that one condition for the retardation of cementite is that the latter must grow under para-equilibrium conditions, i.e. the silicon must be trapped in the cementite. However, this is not a sufficient condition in that it can only be effective in retarding the transformation rate if the overall driving force for the reaction is not large. It is demonstrated that the experience that silicon retards the tempering of martensite requires the presence of lattice defects which can reduce the amount of carbon available for precipitation and the associated driving force.
Article
Full-text available
Galvanizability of high-strength steels was required to improve the fuel efficiency and reduce the materials costs in automotive applications. The standard CMnSi transformation induced plasticity (TRIP) compositions for cold-rolled intercritically processed TRIP grades were shown to be incompatible with the currently used galvanizing processes. No islandlike features were revealed from the scanning electron microscopy (SEM) images of the surface of the steel annealed in the high dew point atmosphere. The temperature of the isothermal holding was found to be too high to obtain a sufficient amount of retained austenite in the microstructure.
Article
A new type of high strength, high toughness, martensitic steel, based on a newly proposed Quench and Partitioning (Q&P) process, is presented. This high strength martensitic grade is produced by the controlled low temperature partitioning of carbon from as-quenched martensite laths to retained inter-lath austenite under conditions where both low temperature transition carbide formation and cementite precipitation are suppressed. The contribution focuses on both the current understanding of the fundamental processes involved and includes a discussion of the technical feasibility of large-scale industrial production of these steels as sheet products. The Q&P process, which is carried out on steels with a lean composition, should be implemented easily on some current industrial continuous annealing and galvanizing lines. In addition, martensitic Q&P sheet steel is characterized by very favourable combinations of strength, ductility and toughness, which are particularly relevant for high strength anti-intrusion automotive parts.
Article
Several researchers have been working for the development of advanced high-strength steel (AHSS) grades to ensure improved fuel economy, and increased passenger safety of vehicles. Projections are that the weight percentage of AHSS steel will increase to 35% by 2015, whereas mild steel will decrease from 55% (in 2007) to 29% in bodies and closures of light vehicles. The AHSS grades that are currently being applied or are under increased investigation by steel appliers, include dual-phase (DP), complex phase (CP) and transformation induced plasticity (TRIP) steels referred to as first generation AHSS. The austenitic twinning induced plasticity (TWIP) steels, lightweight steels with induced plasticity (L-IP) and shear band strengthened steels (SIP) are referred to as second generation AHSS. The second-generation AHSS steels clearly exhibit superior mechanical properties, but these austenitic grades are highly alloyed, resulting in a significant cost. Martensite/ferrite and martensite/austenite microstructures are considered for the third-generation AHSS.
Article
This paper presents a detailed characterization of the microstructural development of a new quenching and partitioning (Q&P) steel. Q&P treatments, starting from full austenitization, were applied to the developed steel, leading to microstructures containing volume fractions of retained austenite of up to 0.15. The austenite was distributed as films in between the martensite laths. Analysis demonstrates that, in this material, stabilization of austenite can be achieved at significantly shorter time scales via the Q&P route than is possible via a bainitic isothermal holding. The results showed that the thermal stabilization of austenite during the partitioning step is not necessarily accompanied by a significant expansion of the material. This implies that the process of carbon partitioning from martensite to austenite occurs across low-mobility martensite–austenite interfaces. The amount of martensite formed during the first quench has been quantified. Unlike martensite formed in the final quench, this martensite was found to be tempered during partitioning. Measured volume fractions of retained austenite after different treatments were compared with simulations using model descriptions for carbon partitioning from martensite to austenite. Simulation results confirmed that the carbon partitioning takes place at low-mobility martensite–austenite interfaces.
Article
The kinetics of athermal martensitic transformation have been studied in a high silicon steel (300 M), beginning with samples which were first partially transformed to bainitic ferrite. It is found that the way in which the volume fraction of martensite increases with undercooling below the martensite start temperature is not greatly influenced by the presence of bainitic ferrite, when any carbon enrichment in the residual austenite is taken into account. the martensitic transformation obeys, within the limits of experimental errors, the same law irrespective of the presence or absence of bainitic ferrite prior to transformation. A new relationship, which takes some account of autocatalysis, has been derived to rationalize the athermal kinetics of martensitic reactions and, within the context of certain approximations, is found to be in reasonable agreement with experimental data. The role of chemical composition variations, of the type normally present in commercial steels, seems to be mainly to extend the temperature range over which most of the martensite reaction occurs in the heterogeneous samples, relative to samples which were given a homogenizing heat treatment.
Article
Three stabilization mechanisms—the shortage of nuclei, the partitioning of alloying elements, and the fine grain size—of the remaining metastable austenite in transformation-induced plasticity (TRIP) steels have been studied by choosing a model alloy Fe-0.2C-1.5Mn-1.5Si. An examination of the nucleus density required for an athermal nucleation mechanism indicates that such a mechanism needs a nucleus density as large as 2.5 · 1017 m−3 when the dispersed austenite grain size is down to 1 µm. Whether the random nucleation on various heterogeneities is likely to dominate the reaction kinetics depends on the heterogeneous embryo density. Chemical stabilization due to the enrichment of carbon in the retained austenite is the most important operational mechanism for the austenite retention. Based on the analysis of 57 engineering steels and some systematic experimental results, an exponential equation describing the influence of carbon concentration on the martensite start (M s) temperature has been determined to be M s (K)=273+545.8 · e −1.362w c(mass pct). A function describing the M s temperature and the energy change of the system has been found, which has been used to study the influence of the grain size on the M s temperature. The decrease in the grain size of the dispersed residual austenite gives rise to a significant decrease in the M s temperature when the grain size is as small as 0.1 µm. It is concluded that the influence of the grain size of the retained austenite can become an important factor in decreasing the M s temperature with respect to the TRIP steels.