Phase-Field Modeling of Austenite-to-Ferrite Transformation in Fe-C-Mn Ternary Alloys

doi:10.11900/0412.1961.2016.00468

Acta Metall Sin

2017, Vol. 53

Issue (6): 760-768 DOI: 10.11900/0412.1961.2016.00468

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Phase-Field Modeling of Austenite-to-Ferrite Transformation in Fe-C-Mn Ternary Alloys

Jun ZHANG^1,²,Wenxiong CHEN²,Chengwu ZHENG²(

),Dianzhong LI²

1 School of Chemistry and Materials Science, University of Science and Technology of China, Hefei 230026, China
2 Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, China

Cite this article:

Jun ZHANG,Wenxiong CHEN,Chengwu ZHENG,Dianzhong LI. Phase-Field Modeling of Austenite-to-Ferrite Transformation in Fe-C-Mn Ternary Alloys. Acta Metall Sin, 2017, 53(6): 760-768.

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Abstract

The effect of Mn on the austenite-to-ferrite transformation has been widely studied by both physical models and experiments due to its technological importance for alloy design in steel industries. In recent years, an increasing interest of this issue is moved onto the effect of alloying element on the migrating interface during the austenite-to-ferrite transformation. For ternary Fe-C-Mn alloys, the interfacial condition is more complicated than that of binary Fe-C alloys in view of the large difference in the diffusivity between the interstitial and substitutional alloying elements. Generally speaking, there are two main concepts, i.e. the paraequilibrium model and the local-equilibrium model, which have been proposed to describe the phase transformation kinetics in ternary Fe-C-Mn alloys based on different assumptions about the diffusion of the substitutional elements. And many modeling attempts have been made to study the effect of Mn on the migration kinetics by using these theories. In this work, a multi-phase-field (MPF) model coupling with a Gibbs-energy dissipation model was developed to simulate the isothermal austenite-to-ferrite transformation in ternary Fe-C-Mn alloys. This model has considered the Mn diffusion inside the migrating interface in a physical manner and takes its effect on the transformation kinetics into account. Comparison simulations were made to analyze the difference in the transformation kinetics and ferrite morphologies with and without considering the energy dissipation at the moving interface. It shows that the incomplete transformation phenomenon does occur due to the Mn diffusion inside interface. The modified MPF model was then used to study the effect of Mn contents on the microstructures and kinetics of the phase transformations. It is found that the ferrite growth along the austenite/austenite boundaries is faster than that in the perpendicular direction. This difference is intensified with increasing the Mn concentration, which hence leads to the ferrite morphology changed from elliptical to flat alike. It also produces a slower transformation kinetics and a larger degree of the incomplete transformation when increasing the Mn concentration.

Key words: phase-field method austenite ferrite Gibbs-energy dissipation incomplete transformation

Received: 21 October 2016

Fund: Supported by National Natural Science Foundation of China (Nos.51371169 and 51401214)

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	Jun ZHANG
	Wenxiong CHEN
	Chengwu ZHENG
	Dianzhong LI

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https://www.ams.org.cn/EN/10.11900/0412.1961.2016.00468 OR https://www.ams.org.cn/EN/Y2017/V53/I6/760

Fig.1 Schematic of the chemical potential of Mn inside the austenite/ferrite interface (2ζ is the physical interface thickness, E₀ is the Mn interface binding energy, μMnγandμMnαare the chemical potentials of Mn in austenite and ferrite, respectively)

Fig.2 Gibbs-energy dissipation (ΔG^dis) as a function of Vζ/DintMn during γ→α transformation with the binding energy E₀=10 kJ/mol (V—interface velocity, ζ—half thickness of the physical interface, DintMn—trans-interface diffusivity of Mn)

Table 1 Physical parameters used in the simulations

Fig.3 Single ferrite α growths between austenite grains γ₁ and γ₂ at 973 K with ΔG^dis (a1~a3) and without ΔG^dis (b1~b3) at time t =20 s (a1, b1), t =40 s (a2, b2) and t =60 s (a3, b3)

Fig.4 Kinetics of ferrite α growth under different conditions at 973 K, comparing with the LE and PE predictions^[18]

Fig.5 Variations of total dissipation of the chemical free energy (ΔG^chem) and ΔG^dis with time due to Mn diffusion inside the interface for austenite-to-ferrite transformation at 973 K

Fig.6 Chemical driving forces for the austenite-to-ferrite transformation as a function of temperature (T) in the Fe-0.1C-2.0Mn, Fe-0.1C-1.0Mn and Fe-0.1C-0.5Mn (mass fraction, %) alloys (T_c—triggering temperature of γ→α transformation)

Fig.7 Simulations of temporal evolutions of the microstructures at 1043 K during the austenite-to-ferrite transformation for Fe-0.1C-2.0Mn (a1~a4) , Fe-0.1C-1.0Mn (b1~b4) and Fe-0.1C-0.5Mn (c1~c4)

Fig.8 Simulations of temporal evolutions of the carbon concentration fields at 1043 K during the austenite-to-ferrite transformation for Fe-0.1C-2.0Mn (a1~a4) , Fe-0.1C-1.0Mn (b1~b4) and Fe-0.1C-0.5Mn (c1~c4)

Fig.9 Ferrite fractions as a function of time during isothermal austenite-to-ferrite transformation at 1043 K for different alloys (fαeq—equilibrium volume fraction of ferrite)

Fig.10 Degrees of incomplete transformation at 1043 K as a function of Mn concentration

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