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All issues Volume 117 / No 5 (2020) Metall. Res. Technol., 117 5 (2020) 502 Full HTML
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Metall. Res. Technol.
Volume 117, Number 5, 2020
Article Number 502
Number of page(s) 7
DOI https://doi.org/10.1051/metal/2020047
Published online 17 August 2020

© EDP Sciences, 2020

1 Introduction

For austenite transformation, most works are focused on its decomposition, such as the formation of ferrite, bainite and martensite, due to the significant influence of these transformations on the relevant mechanical properties. However, the formation process of austenite is also important for the evolution of the microstructure, and eventually for the mechanical properties. Ferrite plus pearlite is the characteristic microstructure in as hot-rolled products of advanced high strength steels for automotive applications, such as dual-phase (DP), transformation-induced plasticity (TRIP) assisted, or quenching and partitioning (Q&P) alloys, whose cold-rolled microstructures would transform into austenite partially during intercritical annealing in a continuous annealing process [1–4]. Therefore, the ferrite-austenite and pearlite-austenite transformation associated with cementite dissolution are an inevitable and vital process in microstructural evolution of automotive steels [5,6].

The formation of austenite from pearlite or pearlite plus ferrite under non-isothermal or isothermal condition in conventional steels has been previously studied [6–8]. There are two critical temperature points for the austenite formation, i.e., Ac1 and Ac3, between which austenite and ferrite co-exist either with or without cementite [9]. The temperature for the full austenitizing or intercritical annealing treatment is determined in practice based on Ac1 and Ac3 points, which could be measured using dilatometer by the change of specific volume due to the lattice structure change accompanied with phase transformation [9]. The transformation into austenite during heating is influenced by chemical compositions and initial microstructure of alloys as well as heating rate [9–11].

Nowadays, the next-generation automotive alloy design concepts of δ-TRIP steels, δ-QP steels and medium-Mn high aluminium TRIP steels to obtain high strength and ultra-ductility is associated with high aluminium addition of 2.5∼6 mass% [12–14]. The addition of aluminium would modify the phase fraction of initial microstructures of hot- or cold-rolled products, and drastically influences the thermodynamics and kinetics of austenitic transformation from ferrite plus cementite. Hence, investigations on the effect of aluminium on the austenite formation are highly necessary for the control of microstructure evolution of high aluminium-added steels.

In the present work, the initial microstructure consisting of ferrite and pearlite was analyzed. In order to clarify the microstructure evolution during intercritical annealing in the high aluminium-added steels, the transformation behaviour from ferrite plus cementite into austenite during heating and isothermal process of two alloys with or without aluminium addition has been investigated. Transformation thermodynamics and kinetics were further analyzed by ThermoCalc in combination with DICTRA and experimental efforts.

2 Experimental procedure

A high aluminium steel, which is referred to the addition of a high content of aluminium in the previous δ-TRIP alloy composition system, has been engaged in this study, and it was labelled as Alloy-1 [12,15–17]. In order to investigate the effect of aluminium on the hot-rolled microstructure and the transformation into austenite, an alloy named as Alloy-2 containing similar compositions but without aluminium was produced for comparisons. The alloys were manufactured as 30 kg ingots in dimensions of 100 × 160 × 230 mm3 in a vacuum furnace. The actual chemical compositions of Alloy-1 are Fe-0.35C-0.57Mn-0.44Si-0.51Cr-2.5Al mass%, and chemical compositions of Alloy-2 are Fe-0.33C-0.59Mn-0.45Si-0.5Cr mass%. The ingots were reheated to 1200 °C before forging into a billet with a section dimension of 60 × 60 mm2. The billets were then reheated to 1200 °C and hot-rolled to sheets with the thickness of ∼4 mm at the temperature between 1100 °C and 900 °C followed by air-cooling.

Equilibrium phase diagrams, including the temperatures of Ae1 and Ae3, were calculated using the software of ThermoCalc with TCFE9 database, and only phases of FCC, BCC and cementite were included in the calculation. DICTRA in ThermoCalc was used to simulate the transformation kinetics combined with TCFE9 and MOBFE4 databases. The initial temperature of kinetics simulation for heating process was the temperature for cementite fully dissolution, and initial phases were ferrite and austenite. The initial compositions of ferrite and austenite were equilibrium compositions at this temperature, which was obtained from thermodynamic calculation. The initial length of ferrite and austenite in both alloys were set as 10 μm for simulation simplification, which is referred to the hot-rolled microstructure. The heating rates of 0.1, 1 and 10  °C · s−1 were used. The interface positions at different temperatures and element profiles of carbon and aluminium were obtained.

The phase transformation was monitored by a push-rod Formastor-FII high-speed dilatometer with radio frequency induction heating. The dilatometer samples are cylindrical with a diameter of 3 mm and length of 10 mm. The heating dilatometric curves were measured to determine the start temperature (Ac1), finish temperature (Ac3) and the pearlite dissolution finish temperature (Afc1) of pearlite plus ferrite to austenite transformation. The determined method of temperature points were referred previous research [10,18]. For high aluminium steels, there is an abnormal expansion at the beginning of pearlite to austenite transformation, accompanying with the slope of the dilatation curve increasing. Then, the slope of the dilatation curve starts to decline, as most of the pearlite is dissolved, and this temperature point was identified as Afc1 [18]. In order to investigate the effect of aluminium on the transformation temperature during reheating, the heating experiment with different heating rates was conducted. Both alloys were heated to 500 °C at a rate of 10 °C · s−1, where no transformation is expected, followed by heating rates of 0.05, 0.1, 0.5, 1, 2, 5 and 10 °C · s−1, respectively, to 1150 °C. In order to validate the accuracy of TCFE9 database during equilibrium thermodynamic calculation for the high aluminium-added alloy, the specimens of alloy1 were soaked at 750 °C, 800 °C, 850 °C, 900 °C, 950 °C, 1000 °C, 1050 °C and 1100 °C for 1 hour, respectively, followed by quenching at a cooling rate of 80 °C · s−1 to hold the high-temperature phase fraction into ambient temperature. The intercritical isothermal annealing treatments of Alloy-1 at 780 °C and Alloy-2 at 720 °C for 15 s, 30 s, 60 s, 120 s, 300 s and 600 s, followed by quenching, were conducted to study the effect of aluminium on the isothermal transformation.

The metallography samples were prepared using standard methods and etched in 4 vol.% Nital, and the relevant microstructures were observed by an electron scanning microscope (JEOL JXA-8530F) operated at 20 kV accelerating voltage. The phase constitutions were quantitatively analyzed by using ImageJ software. The mean inter-lamellae λ of pearlite was derived from mean intercept spacing l, measured in random scanning electronic microscope (SEM) fields, according to equation: l = 2λ [19,20].

3 Results and discussions

3.1 Initial hot-rolled microstructure

The microstructures of both alloys in the as hot-rolled condition are a mixture of ferrite and pearlite (Fig. 1). The initial microstructures are discussed here because they can affect the consequent phase transformation during reheating and isothermal holding [21]. Due to the addition of aluminium element, which is a strong ferrite stabilizer, more ferrite (68 ± 1 vol.%) with the finer grain size of 8.3 ± 0.8 μm is obtained in alloy 1. In comparison, only 55 ± 5 vol.% of ferrite with a grain size of 10.4 ± 0.4 μm is obtained in alloy 2 [1,22–24]. Not only the grain size of ferrite and pearlite was refined, but also the inter-lamellae spacing of pearlite becomes refined from 144 ± 10 nm in Alloy-2 to 52 ± 13 nm in Alloy-1 by aluminium addition (Fig. 1). The aluminium raises Ae1 temperature and therefore increases the undercooling of pearlite reaction [19,25]. The inter-lamellae spacing λ was described as: where σαθ is the interfacial energy per unit area, TE is the eutectoid temperature, and ΔHV is the enthalpy change per unit volume [26]. Even though the aluminium addition increased TE and undercooling TE−T at the same time, the value of consequently decreased, inducing the smaller inter-lamellae spacing λ [25]. Furthermore, the lamellae thickening process is suppressed since the aluminium must diffuse away from cementite during transformation due to the negligible solubility of aluminium in cementite [19,27,28].

thumbnail Fig. 1

Scanning electron micrographs of the as hot-rolled alloys: (a) Alloy-1 and (b) Alloy-2.

3.2 Equilibrium phase and isothermal transformation

The aluminium addition in the steel extensively expands the intercritical region with only ferrite and austenite being presented. According to the thermodynamics calculation by ThermoCalc, the austenite starts to form at the temperature of 759 °C (Ae1) and finishes the transformation at 1043 °C (Ae3) for Alloy-1 (Fig. 2a and Tab. 1). The difference between Ae1 and Ae3 is 284 °C. For Alloy-2, without the addition of aluminium, the temperature range of the intercritical region is only 72 °C (Tab. 1), which is 212 °C less than that for Alloy-1 (Fig. 2a, b). The dissolution of cementite happens very quickly in both two alloys at temperatures between Ae1 and Afe1, and the latter is defined as the point for a full dissolution of cementite into austenite. The interval between Ae1 and Afe1 is 10 °C and 11 °C for Alloy-1 and Alloy-2 (Fig. 2), respectively. Previous work have demonstrated that the pearlite-austenite transformation is completed within less than one second at the temperature higher than 800 °C, and the ferrite to austenite transformation kinetics is relying on the heating rate and the holding temperature [10].

In an approximately equilibrium condition of isothermal holding for 1 hour, the austenite during intercritical annealing was frozen into the final microstructure and revealed as an MA (martensite/austenite) phase in the as-quenched samples. The amount of intercritical austenite was measured by image analysis. The actual intercritical austenite fractions are close to those predicted by the equilibrium phase diagram based on thermodynamics using ThermoCalc (Fig. 2a), except the results at 750 °C. The measured Ac1 temperatures of both alloys are several tens of degree lower than those indicated in equilibrium phase diagrams (Tab. 1). This is likely due to the heterogeneous solute distribution in the hot-rolled microstructure. The higher contents of carbon and manganese are enriched in pearlite, which decreases the starting points of austenite formation in this region. At the same time, the interfacial energy of cementite/ferrite stored in the fine lamella accelerates the austenite nucleation, consequently decreases the austenite formation temperature.

The Ac1 and Ac3 temperatures of Alloy-1 were measured as 700 °C and 1049 °C, respectively, at the condition of approximate equilibrium with an extremely low heating rate of 0.05 °C · s−1 (Fig. 3 and Tab. 1). The intercritical region lasted for 349 °C due to aluminium addition. In Alloy-2, austenite started to form at 659 °C and finished at Ac3 temperature of 798 °C, and the ferrite survived for 139 °C. This significantly depends on the aluminium addition, which restricts the formation of austenite and stabilizes the ferrite. The measurements of Ac3 temperatures in both alloys are very close to the temperatures obtained in the thermodynamic equilibrium phase calculation, due to the homogeneous distribution of elements after long-term heat treatments.

The intercritical isothermal process of two alloys was investigated, including the extremely quick pearlite transformation and relatively slower ferrite-austenite transformation. The intercritical annealing was conducted at 780 °C of Alloy-1 and 720 °C of Alloy-2; the chosen temperature was the temperature of a full dissolution of the cementite, determined through metallography observation. The formation of austenite depends on the annealing duration at a constant temperature [10]. The dilatation curves of the samples during intercritical annealing for 600 s are shown in Figure 4a, and the magnified curves (red dash line) at the first 180 s during isothermal is included. The transformation in Alloy-2 was almost finished before 180 s. But for Alloy-1, the transformation lasted for the whole isothermal process, as the strain decreased consistently. The rate of transformation was reduced with extending isothermal time. The ratio of amount for transformation at the preliminary 180 s to subsequently isothermal process is 5.6:1. After intercritical annealing for different times followed by quenching, the volume percentage of ferrite was measured by metallographic and computational method. With increasing isothermal time from 30 s to 120 s, the ferrite percentage of Alloy-2 decreased from 38.6 ± 1.7 vol.% to 27.1 ± 1.5 vol.%, and then almost remained unchanged with further increasing time. The ferrite percentage of Alloy-1 lowered from 37.8 ± 1.1 vol.% at 30 s to 27.7 ± 1.2 vol.% at 120 s, and to 24.5 ± 1.7 vol.% at 600 s (Fig. 4b). The decreasing of ferrite fraction in Alloy-1 lasted with increasing time to 600 s. The reason for the delayed transformation of Alloy-1 is attributed to the fact that the diffusion of aluminium is necessary for austenite formation, which influences the thermodynamic stability of austenite relative to ferrite. This result indicated that the alloys containing aluminium need more time for the stabilization of transformation during intercritical annealing, despite the higher isothermal temperature. Therefore, a longer intercritical isothermal time is needed for aluminium added steel during continuous annealing or galvanising process in experimental and industrial manufacture.

thumbnail Fig. 2

The change in phase fraction with temperature: (a) Alloy-1 in equilibrium combined with measured results, and (b) Alloy-2 in equilibrium (γ, α and θ represent austenite, ferrite and cementite, respectively).
Table 1

Critical temperatures calculated by thermodynamics and measured by dilatometer, in [°C].

thumbnail Fig. 3

The curves of relative change of length as a function of temperature, showing Ac1 and Ac3 temperatures in quasi-equilibrium condition at a heating rate of 0.05 °C · s−1.
thumbnail Fig. 4

Alloy-1 isothermal at 780 °C and Alloy-2 isothermal at 720 °C: (a) the relative change of length of experimental alloys as a function of time during intercritical annealing, included with the locally magnified curves showing the transformation of Alloy-1 almost finished at the first 180 s, (b) volume percentage of ferrite as a function of isothermal times.

3.3 Transformation during heating

Increasing the heating rate from the quasi-equilibrium condition of 0.05 °C · s−1 to typical heating rate in commercial lines of 10 °C · s−1 (Fig. 5), the austenite transformation start temperature is delayed by 22 °C in Alloy-2 and 39 °C in Alloy-1, respectively, where the aluminium addition has played a significant role. The Afc1 of Alloy-1 is elevated for 36 °C with the increasing of heating rate (Fig. 5a), while Alloy-2 does not exhibit an apparent Afc1 point (Fig. 3). In the conventional plain carbon alloys, the rate for pearlite formation or reverse transformation into austenite mainly relies on the carbon diffusion. The solubility of aluminium in cementite relative to ferrite and austenite is thought to be negligible; therefore, the aluminium needs to partition away from the cementite [28]. As the heating rate increases, the delay in transformation in Alloy-1 becomes more severe, since the diffusion of aluminium is necessary for stage 1 before Afc1. The Ac3 temperatures in Alloy-1 are elevated for 82 °C when the heating rate increases to 10 °C · s−1, owing to a slow transformation rate. The addition of aluminium makes the transformation temperature to be more sensitive to the heating rate.

In order to verify the influence of aluminium on the ferrite-to-austenite transformation behaviour during reheating, the austenite/ferrite (γ/α) interface migration and the evolutions of carbon and aluminium profiles were simulated by DICTRA. In this simulation, the initial condition was taken to be a mixture of austenite and ferrite with the compositions being obtained from thermodynamic equilibrium calculation. A one-dimensional growth of interface was simulated in the model. The grain size of austenite and ferrite was referred to the microstructure of hot rolled materials. The transformation was simulated from the initial temperature of 767 °C and 733 °C for Alloy-1 and Alloy-2, respectively. The γ/α interface moves toward ferrite with ferrite-austenite transformation during heating in each alloy (Fig. 6). The interface migration simulation results show that Alloy-1 needs a higher temperature for the migration of γ/α interface to reach the same position. This indicates that the transformation of Alloy-1 is delayed, which is also confirmed by the equilibrium calculation and experimental work. Compared with the interstitial atom of carbon, substitutional element aluminium has a smaller diffusion coefficient, and its redistributions take times, bringing about the delay of transformation (Fig. 7a, b). At the start of the transformation, there is an aluminium depletion at the austenite interface. The aluminium-enriched ferrite interface moves from left to right, and the content of aluminium in ferrite increases as the phase transformation takes place during the reheating process. The reason for carbon enriching at austenite interface at the initial stage during reheating but disappearing at a higher temperature is closely related to the high diffusion coefficient of carbon (Fig. 7a, c). Moreover, the distribution of carbon in austenite of Alloy-1 is more homogeneous (Fig. 7a, c), due to the faster diffusion rate resulting from higher temperature and larger lattice parameters in austenite of the aluminium added steels [18]. Besides, the transformation of aluminium added alloys is delayed by the elevated heating rate (Fig. 6), resulting from the slower diffusion rate of aluminium and the necessary degree of superheat for transformation.

thumbnail Fig. 5

Heating rate sensitivity on austenite transformation during heating in alloys affected by aluminium addition, measured by using dilatometer with a spectrum of heating rate between 0.05 and 10 °C · s−1: (a) Alloy-1 and (b) Alloy-2.
thumbnail Fig. 6

The ferrite-austenite interface migration during heating at the rate of 0.1 °C/s, 1 °C/s and 10 °C/s, simulated by DICTRA combined with the database of TCFE9 and MOBFE4.
thumbnail Fig. 7

The evolution of carbon and aluminium profiles of two alloys in heating rate of 10 °C/s: (a) carbon in Alloy-1, (b) aluminium in Alloy-1, and (c) carbon in Alloy-2.

4 Conclusion

The critical role of aluminium on the hot-rolled microstructure and the austenite formation has been investigated in this paper. Aluminium refines the inter-lamellae structure of pearlite from 144 ± 10 nm to 52 ± 13 nm in hot-rolled condition, increases the ferrite fraction and refines its grain size. Aluminium, as a strong ferrite stabilizer, significantly elevates both Ae1 temperature from 727 °C to 759 °C and Ae3 from 799 °C to 1043 °C of steels and enlarges transformation temperature range between Ae1 and Ae3. The isothermal transformation of aluminium added steels is obviously delayed, as a result of necessary aluminium diffusion for the thermodynamic stability of austenite relative to ferrite. This has an introduction for the experimental and industrial manufacture, i.e., a longer isothermal time is needed for continuous annealing or galvanising process of high aluminium added steel. For the heating process, the transformation for austenite formation is delayed by both aluminium addition and heating rate increasing, and the transformation rate of high aluminium added steel is more sensitive to the heating rate.

Acknowledgments

The research was financially supported by the National Natural Science Foundation of China (Grant No. 51804072), as well as the Fundamental Research Funds for the Central Universities (Grant No. N170203004), and China Postdoctoral Science Foundation (Grant No. 2018M631802).

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Cite this article as: Peng Chen, Jian Fu, Jingkuan Yang, Xiaowu Li, A critical role of aluminium on austenite formation in high aluminium added steels, Metall. Res. Technol. 117, 502 (2020)

All Tables

Table 1

Critical temperatures calculated by thermodynamics and measured by dilatometer, in [°C].

In the text

All Figures

thumbnail Fig. 1

Scanning electron micrographs of the as hot-rolled alloys: (a) Alloy-1 and (b) Alloy-2.
In the text
thumbnail Fig. 2

The change in phase fraction with temperature: (a) Alloy-1 in equilibrium combined with measured results, and (b) Alloy-2 in equilibrium (γ, α and θ represent austenite, ferrite and cementite, respectively).
In the text
thumbnail Fig. 3

The curves of relative change of length as a function of temperature, showing Ac1 and Ac3 temperatures in quasi-equilibrium condition at a heating rate of 0.05 °C · s−1.
In the text
thumbnail Fig. 4

Alloy-1 isothermal at 780 °C and Alloy-2 isothermal at 720 °C: (a) the relative change of length of experimental alloys as a function of time during intercritical annealing, included with the locally magnified curves showing the transformation of Alloy-1 almost finished at the first 180 s, (b) volume percentage of ferrite as a function of isothermal times.
In the text
thumbnail Fig. 5

Heating rate sensitivity on austenite transformation during heating in alloys affected by aluminium addition, measured by using dilatometer with a spectrum of heating rate between 0.05 and 10 °C · s−1: (a) Alloy-1 and (b) Alloy-2.
In the text
thumbnail Fig. 6

The ferrite-austenite interface migration during heating at the rate of 0.1 °C/s, 1 °C/s and 10 °C/s, simulated by DICTRA combined with the database of TCFE9 and MOBFE4.
In the text
thumbnail Fig. 7

The evolution of carbon and aluminium profiles of two alloys in heating rate of 10 °C/s: (a) carbon in Alloy-1, (b) aluminium in Alloy-1, and (c) carbon in Alloy-2.
In the text

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