trip steels: advanced high strength multiphase steels

14th International Scientific Conference CO-MAT-TECH 2006

CO-MAT-TECH 2006

TRNAVA, 19 - 20 October 2006

TRIP STEELS: ADVANCED HIGH STRENGTH MULTIPHASE STEELS FOR AUTOMOTIVE APPLICATIONS Daniel KRIZAN Research and Development Department, Business Area Cold Rolling and Coating, voestalpine Stahl GmbH, voestalpine-Strasse 3, 4031 Linz, Austria, [email protected] Work performed at Laboratory for Iron and Steelmaking, Ghent University, Technologiepark 903, B-9052 Ghent, Belgium

1 INTRODUCTION Various types of advanced high strength steels (AHSS) have been developed especially taking in mind the weight reduction of passenger cars. The increase in strength of the more conventional HSS, e.g. precipitated strengthened high strength low-alloyed (HSLA) steel, is usually accompanied by a linear decrease in elongation properties. AHSS possess an improved strength without much loss in elongation or in sheet deep drawing potential [1] (Figure 1).

Figure 1: Elongation vs. tensile strength diagram for a variety of sheet steel materials. Note the clear advantage of the DP and TRIP steels, which combine higher formability and high strength. 1

20-21 October 2006, Trnava, Slovak Republic

AHSS are characterized by superior strain hardening properties and large uniform elongations. Both properties make these steels highly attractive for the automotive industry, in particular in stretch forming applications and in high-energy impact situations, where the energy absorbing capacity is particularly important for safety considerations. These steels include dual phase (DP), transformation induced plasticity (TRIP), deformation-induced twinning steels (TWIP), complex phase (CP) and martensitic steels [2]. TRIP steels are ideally suited for safety-related structural parts for automotive applications due to their superior high strain rate performance, which results in a large dynamic energy absorption. Low alloyed TRIP steels have a multiphase microstructure, consisting of two ferritic micro-constituents, ferrite (~60 vol. %) and bainite (~25 vol. %), and a smaller volume fraction of high C retained austenite (~15 vol. %) as can be seen in Figure 2.

γret α

αB

Figure 2: Typical microstructure of a cold rolled intercritically annealed TRIP steel; ferrite –gray (α), bainite –dark (αB), retained austenite –white (γret). Although the retained austenite is the minor phase present in the TRIP microstructure, it has a major effect on the mechanical properties of TRIP steels. Si, Al, P additions alone or in combination suppress the carbide formation during the isothermal bainitic transformation. This allows the remaining austenite to be enriched in C, which results in its room temperature stabilization. The retained austenite transforms to martensite during straining of TRIP steels, resulting in an increased work hardening rate at higher strain. This causes the superior strength-ductility balance of TRIP steels shown in Figure 1. The goal of the present paper is to give in-depth explanation on the basic physical metallurgical aspects related to TRIP steels, such as TRIP effect and retained austenite stability. Based on these fundamentals, the chemical composition, mechanical properties and processing of the current cold rolled intercritically annealed TRIP steels are outlined.

2 TRANSFORMATION INDUCED PLASTICITY The first observation of an unexpected increase in formability, due to the austenite to martensite transformation in Fe-Ni alloys, was observed in 1937 by Wassermann [3]. In 1967 Zackay et al. [4] described how the transformation in highly-alloyed homogeneous metastable austenitic steels was indeed the reason for the ductility enhancement. The effect, TRansformation Induced Plasticity, was abbreviated as TRIP. Researchers at Nippon Steel Corporation showed that austenite stabilization also occurred during an isothermal bainitic 2


transformation, a process often referred as “austempering”, which followed an intercritical α+γ annealing of low alloy Si-bearing medium-C (0.12 - 0.55 %) CMn steel [5]. In this new class of low-alloyed TRIP steels the austenite is present as a disperse phase [6]. In the published literature different explanations were put forward to explain the TRIP effect in low-alloyed TRIP steels [7-9]. The most accepted one is given by Yokoi et al. [7], asserting that since the C hardening is much higher for the martensite than for the austenite phase and the volume expansion due to this transformation results in plastic deformation and work hardening of the surrounding ferrite, a localized strengthening is obtained. These effects postpone further deformation in this area and move the martensitic transformation to neighbouring areas (see Figure 3), leading to a delay in the onset of macroscopic necking and consequently, to higher values of uniform and total elongation.

retained austenite

γret

Martensite α’

•Ms-M d adjusted to have Msσ~RT •Strain induced transformation: - Uniform elongation higher - Increased strength

Figure 3: Schematic explanation of the TRIP effect showing that the retained austenite transforms to martensite when the stress is applied and the martensitic transformation occurs in neibouring areas in sequence to disperse strain [7].

Two factors determine the TRIP effect, causing enhancement of the mechanical properties of TRIP steels: the volume fraction and the stability of the retained austenite. The optimal volume fraction of the retained austenite for a significant TRIP effect to occur is reported to be in the range of 10-20 vol. % [10]. Moreover, the volume fraction of the retained austenite directly determines the C content and grain size of the retained austenite, its two main stabilization factors. The stability of the retained austenite dictates when the strain-induced martensitic transformation (SIMT) will be triggered during straining of TRIP high strength steel (Figure 4). Unstable retained austenite transforms almost immediately upon deformation, increasing work hardening rate and formability during the stamping process. At the appropriate stability of the retained austenite, the SIMT begins only at strain levels beyond those produced during stamping and forming, and the retained austenite is still present in the final part and it can transform to martensite in the event of a crash, providing greater crash energy absorption. If the retained austenite is too stable, the SIMT may start beyond the uniform elongation. In this case, no additional work hardening is expected and this delayed TRIP effect will not contribute to the enhancement of mechanical properties. 3

Stress


γret unstable

γret too stable

no TRIP effect

no TRIP effect

appropriate γret stability pronounced TRIP effect

Strain Figure 4: Schematic representation of the optimal area for the SIMT associated with a pronounced TRIP effect.

3 RETAINED AUSTENITE STABILITY As already mentioned in the previous paragraph, one of the key factors to make use of the TRIP effect is the control of the SIMT during straining. The temperature dependence of the austenite to martensite transformation is schematically illustrated in Figure 5. Below the Ms temperature, the chemical driving force is large and austenite readily transforms to martensite without deformation. At temperatures above the Ms temperature, the metastable austenite can transform to martensite if a mechanical driving force is added to the chemical driving force. In the temperature range between the Ms temperature and the Msσ temperature, martensite nuclei can be activated by elastic stresses, which rapidly triggers the austenite to martensite transformation at the start of the straining. At the Msσ temperature, the mode of transformation changes from stress-assisted to strain-induced. At this temperature, the stress needed to initiate the martensitic transformation is equal to the yield strength of the austenite. Above the Msσ temperature, yielding of the austenite is initiated by plastic deformation and the martensite nucleation is controlled by plastic strain. Olson et al. [11] assumed that the shear band intersections are the nucleation sites for the SIMT as the shear band intersection sites are areas of high stress concentration. Above the Md temperature, austenite is stable even when plastically deformed, and no transformation occurs. The stability of retained austenite has to be adjusted so that room temperature is above the Msσ temperature and below the Md temperature to induce a pronounced TRIP effect.

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14th International Scientific Conference CO-MAT-TECH 2006 Stress-assisted nucleation (yielding by transformation)

Strain-induced nucleation

σ

σ Yield strength of

YS

γ

(initial yielding by slip)

α

α

α

σ τ α’ γ

γ α’ α

B

B

σ

σ

Athermal

No

martensite

martensite σ

M Md ). The definition of the Msσ temperature and the location of room temperature (RT) to ensure the SIMT at this temperature is also illustrated.

4 STABILIZATION FACTORS The stability of the retained austenite depends on a number of stabilization factors [12]. Three main stabilization factors are listed below: • The C content of the retained austenite: The C content lowers the Ms temperature and therefore stabilizes retained austenite. The optimal C content of the retained austenite in TRIP steel has to be in the range of 0.5-1.8 wt. % in order to provide the desirable TRIP effect [13]; • The size of the retained austenite islands: The Ms temperature of most TRIP steels is above room temperature. Smaller retained austenite particles contain less potential nucleation sites for transformation to martensite and consequently require a greater total driving force for the nucleation of martensite. This will lower the Ms temperature to below room temperature. It has been suggested that the grain size of the retained austenite must be in the range of 0.01 µm to 1 µm to ensure the TRIP effect in lowalloyed multiphase TRIP steels. Larger retained austenite particles may already be partially transformed to martensite or transform to the martensite at the early stages of straining. Particles smaller than 0.01 µm do not undergo the SIMT [13]; • The volume fraction of the retained austenite: As mentioned previously, the optimal volume fraction of the retained austenite for a pronounced TRIP effect is reported to be in the range of 10 - 20 vol. % [10]. Smaller amounts of retained austenite cannot ensure a significant TRIP effect, since the C content may then be too high to result in the SIMT. Large amounts of retained austenite have a low C content, leading to a low stability of the retained austenite.

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5 CHEMICAL COMPOSITION AND MECHANICAL PROPERTIES Current low alloyed TRIP steels are characterized by a very low content of alloying elements of about 3.5 wt. %. Conventional TRIP steel compositions are usually based on the original 0.12-0.55 wt. % C, 0.2-2.5 wt. % Mn, 0.4-1.8 wt. % Si concept proposed by Matsumura et al. [5]. Table 1 reviews some of the recent alternative chemical compositions for cold rolled TRIP steels. Table 1: Alternative compositions for cold rolled TRIP steels (in wt. %) Reference C Mn Al Si De Meyer [14] 0.19 1.57 1.46 De Meyer [14] 0.31 1.57 1.23 0.34 Barbe et al. [15] 0.24 1.66 0.58 0.42 Krizan [1] 0.23 1.65 1.14 0.45

P 0.073 0.076

Ti 0.085

The C content plays a key role in the composition. Its distribution between the main microstructural constituents is fundamental to the properties of the material: it should be enriched as much as possible in the retained austenite in order to have the Msσ temperature of the retained austenite 15 °C-25 °C below RT to obtain the best mechanical properties. Whereas original laboratory TRIP steels could have a C content as high as 0.4 wt. %, current TRIP steels contain typically 0.20 - 0.25 wt. % C or less for reasons of weldability. The typical Mn content in low alloy TRIP steel is ~ 1.5 wt. % Mn to achieve desirable ferrite/austenite volume fraction during intercritical annealing and required hardenability. High Mn contents (~ 2.5 wt. %) are not favorable as they lead to banding in the microstructure and excessively stabilized retained austenite [6]. Si, besides being a strong ferrite stabilizer and an effective solid solution strengthener, inhibits the formation of cementite during the austempering stage. This allows an additional C-enrichment of remaining austenite in this stage, resulting in its subsequent stabilization at RT. Further studies have however revealed important difficulties in hot dip galvanizing of high Si (~1.5 wt.%) cold rolled TRIP steels [16]. The high Si content results in film-forming surface oxides which prevent the formation of the inhibition layer during hot dip galvanizing, causing a pour wettability of TRIP thin sheet steel. This is the main reason to keep the Si content of TRIP steels low. This can be achieved by a partial or full substitution of Si by elements with similar properties [17]. Al [14] and P [15] are the elements of choice in this matter. Both, in addition to being strong ferrite stabilizers, effectively inhibit the formation of cementite during austempering similar to Si. Although low-Si and even Si-free compositions have been proposed, there is at least one very good reason not to remove Si altogether and have at least 0.3 - 0.8 wt. % Si in TRIP steels. Si seems to prevent the most effectively the formation of cementite during the austempering stage. Ideally there should therefore only be a partial replacement of ~ 1 wt. % Si by ~ 1 wt. % Al [14]. By employing the above-mentioned alloying concepts tensile strength from 600 MPa up to 900 MPa and total elongation in the range of 17 % - 36 % can be achieved (Figure 6). The current development of TRIP steels focuses on steels with strength level achieving 1 GPa at appropriate formability properties, i.e. the total elongation of > 18 %. One way to meet these requirements is to increase the C content. This is however directly associated with a deterioration of weldability. Another approach to increase the strength of TRIP steels without deteriorating their weldability by an increase of the C content, is the addition of microalloying elements, such as Ti, Nb and V [1, 18, 19] . Figure 6 shows how the target for the mechanical properties can be fulfilled by using 0.085 wt. % Ti in the cold rolled CMnAlSiP TRIP steel.

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1050

1050

1000

1000 950

Si

900 850 800

Si+P

Si+Al

P

750 Al

700 650 600

Al+P

550 500 450 400

Yield Strength, Tensile Strength, MPa

Yield Strength, Tensile Strength, MPa

950

900 850 800 750 700 650 600 550 500 450 400

350

350

10 12 14 16 18 20 22 24 26 28 30 32 34 36 38 40

10 12 14 16 18 20 22 24 26 28 30 32 34 36 38 40

Total elongation, %

Total elongation, %

Figure 6: Yield strength and tensile strength vs. total elongation for different types of TRIP steels (left) and a Ti micro-alloyed TRIP steel (right; open circles: yield strength; open squares: tensile strength) [20].

6 PROCESSING The processing of cold rolled intercritically annealed TRIP steels consists of the following stages: reheating, hot rolling and coiling, cold rolling and final thermal treatment. Present paragraph focuses on the thermal treatment - annealing of cold rolled TRIP steels, in which the final microstructure and resulting mechanical properties are precisely adjusted. Cold rolled intercritically annealed TRIP steels are usually processed using a thermal cycle such as the one shown schematically in Figure 7. It consists of five distinct stages: heating, intercritical annealing, rapid cooling, austempering and final cooling. The initial microstructure of TRIP steel after coiling at usually 600 °C and subsequent cold rolling consists of cold rolled ferrite and pearlite, latter containing 0.8 wt. % C. During the initial rapid heating of annealing step, the cold rolled ferrite recrystallizes and the cementite starts to dissolve. Above the Ac1 temperature, the pearlite starts transforming to the austenite. During intercritical annealing, the austenitic transformation continues so that the C content of the intercritical austenite gradually decreases to approximately 0.3-0.4 wt. %. The temperature and time are adjusted so that approximately 50 vol. % of the ferrite and 50 vol. % of the intercritical austenite is present in the microstructure at the end of intercritical annealing. After intercritical annealing the steel is cooled to the isothermal bainitic transformation temperature. The cooling rate between intercritical annealing and isothermal bainitic transformation stage, i.e. austempering, must be higher than 10 °C/s in order to avoid the formation of “new” pro-eutectoid ferrite and pearlitic transformation. The subsequent austempering results in a further C-enrichment of the remaining austenite to about 1.5 wt . %. In the final, relatively slow, cooling stage some of the austenite may transform to martensite, especially if the Al content of TRIP steel is high [6]. The final microstructure of cold rolled intercritically annealed TRIP steel after entire processing consists of ferrite, bainite, retained austenite and in some cases a small amount of athermal martensite.

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2. Intercritical annealing: - Form γ with sufficient hardenability

Temperature

time: 2-4 min temperature: 750-800°C

1. Heating: -Recrystallisation of α -Dissolution of cementite -Formation of γ (T>Ac1)

~ -5 C/s 3. Rapid cooling: - Avoid formation of « new» α - Avoid formation of pearlite

~ -10-50 C/s 4. Isothermal Bainitic Transformation: - Enrichment of retained γ with carbon

time: 4-8 min temperature: 350-490°C ~ +5-20 C/s

~ -1-10 C/s 5. Final cooling: -Martensite formation (Ms>RT)

Time Figure 7: Schematic of the thermal treatment for cold rolled low alloy TRIP steel: the main features of the five processing stages are indicated [6].

7 CONCLUSIONS Cold rolled intercritically annealed multiphase TRIP sheet steels can now successfully be produced on an industrial scale. TRIP steels with tensile strengths in the range of 600 MPa800 MPa have been extensively tested and are currently being used by the automotive industry for B-pillars and A-pillars. The micro-alloyed TRIP steels are currently under industrial development and will become available for car producers in the near future. New alternative heat treatment strategies with the chemical compositions very close to that one used for TRIP steels are recently extending the segment of the AHSS for automotive applications. Higher bainite fraction TRIP steels (HBF-TRIP) show an excellent strengthstretch-flangeability balance [21], greatly welcomed by car manufacturers. Furthermore, the quenching and partitioning process [22, 23], which make use of C partitioning between quenched martensite and retained austenite, in the absence of carbide formation, extends the strength of current TRIP steels. Based on this can be stated that the family of AHSS with a TRIP-like behavior offers nowadays a challenging scientific playground with sound industrial scope. Acknowledgements The author wish to thank his former colleagues from Laboratory for Iron and Steelmaking for many fruitful discussions regarding TRIP steels. Financial assistance of OCAS NV, ARCELOR Research Centre, is also gratefully acknowledged.

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[19]

[20]

[21]

[22]

[23]

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Reviewer Dr. Andreas Pichler, voestalpine Stahl GmbH, voestalpine-Strasse 3, A-4031 Linz, Austria

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