Transformation-induced plasticity in multiphase steels subjected to thermomechanical loading.

HAL Id: hal-00513963

https://hal.archives-ouvertes.fr/hal-00513963

Submitted on 1 Sep 2010

HAL is a multi-disciplinary open access archive for the deposit and dissemination of sci- entific research documents, whether they are pub- lished or not. The documents may come from teaching and research institutions in France or

L’archive ouverte pluridisciplinaire HAL, est destinée au dépôt et à la diffusion de documents scientifiques de niveau recherche, publiés ou non, émanant des établissements d’enseignement et de recherche français ou étrangers, des laboratoires

subjected to thermomechanical loading.

Denny Tjahjanto, Sergio Turteltaub, Akke S.J. Suiker, Sybrand van der Zwaag

To cite this version:

Denny Tjahjanto, Sergio Turteltaub, Akke S.J. Suiker, Sybrand van der Zwaag. Transformation- induced plasticity in multiphase steels subjected to thermomechanical loading.. Philosophical Maga- zine, Taylor & Francis, 2009, 88 (28-29), pp.3369-3387. �10.1080/14786430802438150�. �hal-00513963�

For Peer Review Only

Transformation-induced plasticity in multiphase steels subjected to thermomechanical loading.

Journal: Philosophical Magazine & Philosophical Magazine Letters Manuscript ID: TPHM-08-Feb-0051.R1

Journal Selection: Philosophical Magazine Date Submitted by the

Author: 04-Jul-2008

Complete List of Authors: Tjahjanto, Denny; Delft University of Technology, Aerospace Engineering

Turteltaub, Sergio; Delft University of Technology, Aerospace Engineering

Suiker, Akke; Delft University of Technology, Aerospace Engineering

van der Zwaag, Sybrand; Delft University of Technology, Aerospace Engineering

Keywords: computational mechanics, plasticity of metals Keywords (user supplied): transformation-induced plasticity, Multiphase steels, Thermomechanics

Note: The following files were submitted by the author for peer review, but cannot be converted to PDF. You must view these files (e.g. movies) online.

Tjahjanto_etal-PM2008-v2.tex BibDDT.bib

For Peer Review Only

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55

For Peer Review Only

Philosophical Magazine

Vol. 00, No. 00, Month 2008, 1–19

RESEARCH ARTICLE

Transformation-induced plasticity in multiphase steels subjected to thermomechanical loading

D.D. Tjahjanto^†, S. Turteltaub^∗, A.S.J. Suiker and S. van der Zwaag Faculty of Aerospace Engineering, Delft University of Technology

Kluyverweg 1, 2629 HS Delft, the Netherlands

{Email: S.R.Turteltaub,A.S.J.Suiker,S.vanderZwaag}@tudelft.nl

(v3.1 released April 2006)

The behavior of transformation-induced plasticity steels subjected to combined thermome- chanical loading is studied at the microscale by means of numerical simulations. The mi- crostructure is composed of an austenitic phase that may deform plastically and/or transform into martensite, and a ferritic phase that may deform plastically. The micromechanical mod- els capturing these effects are derived from a thermodynamical framework, which has been extended from previous works in order to adequately account for the thermal contributions to the kinematics and the Helmholtz energy. The models are used in numerical simulations on a polycrystalline sample composed of an aggregate of multiple austenitic and ferritic grains of various orientations. The thermomechanical response of the sample is studied under (i) isothermal straining at different temperatures above the martensitic start temperature, and under (ii) different paths of straining and cooling to temperatures below the martensitic start temperature. The first type of analysis shows that at lower temperatures the transformation mechanism is more dominant than the plasticity mechanism, whereas the converse occurs at higher temperatures. The second type of analysis illustrates that, in comparison to a bench- mark initially stress-free sample at room temperature, the transformation rate under straining is higher when performed on an pre-cooled sample, but that the transformation rate under cooling is lower when carried out on a pre-strained sample. The results of this analysis indi- cate that, for optimizing the formability of this class of steels, it is recommendable to make a judicious choice regarding the thermomechanical loading parameters during manufacturing processes.

1. Introduction

Multiphase steels assisted by transformation-induced plasticity (i.e., TRIP steels) have a superior combination of strength and ductility characteristics [1]. This is because at the microscale the grains of retained austenite in a TRIP steel mi- crostructure may transform into martensite upon mechanical and/or thermal load- ing, thereby inducing plasticity in the surrounding phases, i.e., a TRIP effect is generated. In order to accurately predict the overall behavior of this class of steels during manufacturing and/or operation, a thorough knowledge of the evolution of TRIP steel microstructures under thermomechanical loading is necessary.

In the last decades, various micromechanically-based models for martensitic transformations have been developed, see e.g., [2–18]. However, these models mainly focus upon mechanically-driven martensitic transformations, without tak- ing into account the effect of temperature variations on the transformation behav- ior. On the other hand, the characteristics of martensitic transformations in TRIP

†Current address: Max-Planck-Institut f¨ur Eisenforschung GmbH, Max-Planck-Str. 1, 40237 D¨usseldorf, Germany, [email protected]

∗Corresponding author

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

For Peer Review Only

steel microstructures during cooling have been investigated in recent experimen- tal and modeling works [19–24]. Several empirically- and thermodynamically-based models have been proposed for predicting the transformation behavior in multi- phase TRIP steels subjected to cooling. The stability of retained austenite in these models is characterized by the martensitic start temperature,M_s, which is a func- tion of microstructural parameters, such as the carbon concentration and the grain size. Typically, the expression for theM_stemperature is based on the classical em- pirical formulation of Andrews [25]. Nevertheless, the application of such models is limited to thermally-driven martensitic transformations.

In the present work, the thermomechanical model originally proposed by Turteltaub and Suiker [26–28] and further refined by Tjahjantoet al. [29] is adapted to study the behavior of TRIP steels under thermomechanical loadings. In previous works this model has been used to study the effect of microstructural parameters on the behavior of TRIP steels [30, 31]. A modification of the model is required to ac- count for previously neglected thermal effects on the kinematics and the Helmholtz energy. The incorporation of the thermal expansion/contraction effect in the de- formation is derived in the spirit of a model for shape-memory alloys proposed by Anand and Gurtin [32]. Simulations are performed to analyze the effects of microstructural parameters on the transformation and plasticity behavior upon combinations of thermal and mechanical loadings. The outcome of the analyses provides a deeper insight into the TRIP effect in multiphase carbon steels sub- jected to thermomechanical loading. The occurrence of local, crystalline damage in the relatively brittle martensitic product phase is not included in the present model.

This aspect has been discussed elsewhere [33, 34] for TRIP steel microstructures subjected to mechanical loading at ambient temperature.

This paper is organized as follows: The models for the thermomechanical behavior of the austenitic and ferritic phases in a TRIP steel are presented in Sections 2 and 3, respectively. The models are subsequently used in numerical simulations on TRIP steel microstructures subjected to thermomechanical loadings, as discussed in Section 4. Here, for simplicity, the differences between the ferritic and bainitic phases are ignored in the TRIP steel microstructure, and pre-existing martensite generated during the manufacturing process is not included. Concluding remarks on the simulation results are given in Section 5.

2. Thermo-elasto-plastic-transformation model for single-crystalline austenite

2.1. Kinematics, deformation gradients and entropy densities

In order to quantify the different mechanisms characterizing the response of a multiphase steel, the total deformation gradientF and the entropy densityη (per unit mass) in a material point are decomposed as follows (see [29]):

F =F_eF_pF_tr and η=η_e+η_p+η_tr, (1) where F_e, F_p and F_tr are, respectively, the elastic, plastic and transformation deformation gradients, see Figure 1, andη_e,η_p andη_trare the conservative, plastic and transformation parts of the entropy density, respectively. Each material point is meant to represent an infinitesimal neighborhood that may contain a mixture of austenite and one or more crystallographically distinct arrangements of martensite (referred to as transformation systems, see Figure 1). The volume fractions of the transformation systems α (= 1, . . . , M) at a material point x in the reference 1

2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

For Peer Review Only

X

Fp

Fe

Reference configuration

Current configuration Intermediate I

configuration

F(x) Infinitesimal

neighborhood of x

y(x) m_A

n_A

Ftr

Intermediate II configuration Martensitic platelet

(transformation system)

d b x

Slip system in FCC austenite

Figure 1. Schematic representation of the decomposition of the deformation gradientF into elastic-, plastic- and transformation parts and the corresponding configurations.

configuration and at time t are denoted as ξ^(α) = ˆξ^(α)(x, t), and are such that 0≤ξ^(α) ≤1 and 0≤P_M

α=1ξ^(α) ≤1. The volume fraction of the (untransformed) austenite is denoted as ξ_A, and is defined as ξ_A = 1−P_M

α=1ξ^(α). In accordance with the transformation model developed in [26–28] (see also [29]), the time rate of change of the transformation deformation gradient and the transformation part of entropy density are, respectively, given by

F˙_tr= XM

α=1

ξ˙^(α)b^(α)⊗d^(α) and η˙_tr= XM

α=1

ξ˙^(α)λ^(α)_T

θ_T , (2)

where a superimposed dot denotes a material time derivative. The vectorsb^(α) and d^(α) are, respectively, the transformation shape strain vector and the normal to the habit plane of a transformation systemα(measured in the reference configura- tion),λ^(α)_T is the latent heat of a transformation systemα andθ_T is the theoretical transformation temperature (i.e., the temperature at which austenite can trans- form isothermically into a specific system α of martensite at zero stress, without dissipation, and in the absence of an internal barrier). In rate form, the plastic deformation gradient,F_p, and the plastic entropy density,η_p, are given by

L_p= ˙F_pF_p⁻¹ =

NA

X

i=1

˙

γ⁽ⁱ⁾m⁽ⁱ⁾_A ⊗n⁽ⁱ⁾_A and η˙_p=J_tr

NA

X

i=1

˙

γ⁽ⁱ⁾φ⁽ⁱ⁾_A , (3)

where ˙γ⁽ⁱ⁾ =ξ_Aγ˙_A⁽ⁱ⁾/J_tr represents the “effective” plastic slip rate of slip systemi (= 1, . . . , N_A), with ˙γ_A⁽ⁱ⁾ the plastic slip rate in the untransformed austenite and J_tr = detF_tr. The previous expressions are based on the assumption that plastic deformations evolve only in the (untransformed) austenitic region. Plastic deforma- tions that occurred in the martensitic region prior to transformation are assumed to be inherited, but further deformation in the martensitic phase is modeled as elastic, in accordance with experimental observations reported in [35]. In (3), the vectors m⁽ⁱ⁾_A and n⁽ⁱ⁾_A represent, respectively, the slip direction and the slip plane normal of slip system i. In the second intermediate configuration, both vectors are orthogonal, unit vectors. The parameterφ⁽ⁱ⁾_A represents the change in entropy density per unit slip along slip systemiin the austenite.

The present model does not explicitly resolve the kinematics and kinetics at the length scale of individual dislocations (see [18] for such sub-micron model). In order to take into account the local elastic strains of the austenitic lattice associated 1

2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

For Peer Review Only

with the presence of dislocations, a strain-like scalar variable β is introduced as:

β:=b√

ρ_d, wherebis the magnitude of Burger’s vector and ρ_d measures the total dislocation line per unit volume (see also [36]). The rate of change of the elastic strain-like variable, ˙β, is taken to be linearly dependent of the rate of change of the plastic slip, ˙γ⁽ⁱ⁾, as follows:

β˙=

NA

X

i=1

w⁽ⁱ⁾γ˙⁽ⁱ⁾= ξ_A J_tr

NA

X

i=1

w⁽ⁱ⁾γ˙_A⁽ⁱ⁾, (4)

where the functions w⁽ⁱ⁾ depend on the slip resistance, as will be discussed in Section 2.3.

2.2. Helmholtz energy and constitutive relations

The elastic deformation gradient F_e that incorporates the thermal expansion (or contraction) under a temperature change fromθ₀ toθ can be expressed as

F_e=F_e^?(I +A(θ−θ₀)) , (5) where F_e^? represents the isothermal elastic deformation gradient (measured at a temperatureθ), while the termI+A(θ−θ₀) accounts for the stretch due to thermal expansion (or contraction). Here,I defines the second-order identity tensor andA is the effective thermal expansion tensor. The tensorA is obtained as the volume average of the thermal expansion tensor in the austenite and martensite in the second intermediate configuration, i.e.,

A= 1 J_tr

Ã

ξ_AA_A+ (1 +δ_T) XM

α=1

ξ^(α)A^(α)

!

, (6)

withA_AandA^(α) the thermal expansion tensors of the austenite and the marten- sitic transformation systemα, respectively, andδ_T=b^(α)·d^(α)the change in volume due to transformation. The thermal expansion tensors are assumed to be isotropic, and given byA_A=α_AI andA^(α)=α_MI, whereα_Aand α_M are, respectively, the coefficients of thermal expansion of the austenite and martensite.

In accordance with the formulation presented in [29], the Helmholtz energy den- sityψ(per unit mass) is taken to be a function of the state variablesF_e^?, θ, βandξ, whereθis the temperature and ξ={ξ^(α)|α= 1, . . . , M} is a vector whose entries are the volume fractions of the martensitic transformation systems. Further,ψ is written as the sum of the (elastic) strain energy density ψ_m, the thermal energy densityψ_th, the defect energy density ψ_d, and the surface energy densityψ_s, i.e.,

ψ=ψ_m+ψ_th+ψ_d+ψ_s. (7) The elastic strain energyψ_mis defined as a quadratic function of the elastic strain.

Moreover, since the (unconstrained) thermal stretch does not contribute to the elastic strain energy density ψ_m at a temperature θ, the elastic strain energy is given by

ψ_m= 1

2ρ₀J_trCE_e^?·E_e^?, where E_e^? = 1 2

¡(F_e^?)^TF_e^?−I¢

. (8)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

For Peer Review Only

In the above expressions, E_e^? defines the elastic Green-Lagrange strain measured at a temperatureθ, ρ₀ is the mass density in the reference configuration andC is the effective elasticity tensor, which is given by [26, 27]

C= 1 J_tr

Ã

ξ_ACA+ (1 +δ_T) XM

α=1

ξ^(α)C^(α)

!

, (9)

where CA and C^(α) are, respectively, the elasticity tensors of the austenite and the martensitic transformation systemα. The tensorCA is characterized by three independent elastic moduli,κ^A_i (i= 1,2,3) and the tensorsC^(α) are determined by six independent elastic moduli, κ^M_i (i = 1, . . . ,6). In addition, the elastic stiffness tensors of the austenite and martensite depend on their crystallographic orienta- tions, see [26] for details, and may be functions of temperature; however, in the present model the temperature dependency is neglected for reasons of simplicity.

For metals, the magnitude of the thermal expansion tensorAis much smaller than E_e^?, hence quadratic terms in A can be ignored. Consequently, the elastic strain tensorE_e^? can be written in terms of the overall elastic Green-Lagrange strainE_e and the thermal expansion tensorAas follows:

E_e^?=E_e−B(θ−θ₀) , where B = 1 2

¡(F_e^?)^TF_e^?A+A(F_e^?)^TF_e^?¢

. (10) In view of (10)₂, B can be interpreted as the thermal expansion tensor with a correction for the elastic stretch. Observe that if the elastic deformation is relatively small, thenB ≈A. From (8)₁ and (10)₁, the elastic strain energy densityψ_mcan be expressed in terms of the overall elastic strainE_e, the temperature θ and the martensitic volume fractionsξ, i.e.,

ψ_m(E_e, θ,ξ) = 1

2ρ₀J_tr(ξ)C(ξ)E_e·(E_e−2B(ξ) (θ−θ₀)) . (11) Note that in the above expression for the elastic strain energy density ψ_m the higher-order terms in B are neglected, as well as the implicit dependency of B uponE_e and θ.

The expressions for the thermal energyψ_th, the defect energyψ_d and the surface energyψ_s are given by (see [29]),

ψ_th(θ,ξ) =h(ξ) µ

θ−θ_T−θln µ θ

θ_T

¶¶

−η_Tθ+ XM

α=1

ξ^(α)λ^(α)_T , (12)

ψ_d(β,ξ) = 1

2ρ₀ J_tr(ξ)ω_Aµ(ξ)β² and ψ_s(ξ) = χ

`₀ρ₀ XM

α=1

ξ^(α)

³

1−ξ^(α)

´ , (13)

where θ_T is the theoretical transformation temperature, η_T is the conservative entropy density measured at the transformation temperature θ_T, ω_A is a scaling factor for the defect energy in the austenitic phase, χ is the interface energy per unit area and`₀ is a length scale parameter. Furthermore, h(ξ) andµ(ξ) represent the effective specific heat and the effective equivalent shear modulus, which are 1

2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

For Peer Review Only

given by (see [29])

h(ξ) =ξ_Ah_A+ XM

α=1

ξ^(α)h^(α) and µ(ξ) = 1 J_tr

Ã

ξ_Aµ_A+ (1 +δ_T) XM

α=1

ξ^(α)µ^(α)

! , (14) withh_A and h^(α) being the specific heats of the austenite and martensite, respec- tively, andµ_Aandµ^(α)being the equivalent isotropic shear moduli of the austenite and martensite, respectively. The complete expression for the Helmholtz energy density per unit mass follows from substituting (11), (12), (13)₁, and (13)₂ into (7).

As pointed out in [29], the mechanical constitutive relation may be formulated in terms of the second Piola-Kirchhoff stressS, where the expression follows from differentiating the Helmholtz energy with respect to the elastic Green-Lagrange strainE_e:

S = ρ₀ J_tr

∂ψ

∂E_e =C(E_e−B(θ−θ₀)) . (15) In a similar fashion, the thermal constitutive relation in terms of the reversible entropyη_eis obtained by taking the derivative of the Helmholtz energy with respect to the temperatureθ, leading to

η_e=−∂ψ

∂θ =hln µ θ

θ_T

¶

+η_T+J_tr

ρ₀CE_e·B. (16) It is emphasized that, in contrast to the model presented in Tjahjantoet al. [29], the mechanical constitutive relation, (15), and the thermal constitutive relation, (16), are fully coupled.

2.3. Driving forces, kinetic relations and hardening law

In the above section it has been shown that the Helmholtz energy contains four components, see (11), (12), (13)₁ and (13)₂. In correspondence with these four terms, the transformation driving force f^(α) for a martensitic system α may be decomposed as

f^(α)=f_m^(α)+f_th^(α)+f_d^(α)+f_s^(α), (17) where fm^(α), f_th^(α), f_d^(α) and fs^(α) represent the (bulk) mechanical contribution, the thermal contribution, the defect energy contribution and the surface energy contri- bution to the transformation driving force, respectively. Using the procedure given in [26, 29], the bulk mechanical part of the transformation driving force can be obtained from the resolved stress contribution and the derivative of the Helmholtz energy with respect to the martensitic volume fraction,ξ^(α), as

f_m^(α)=J_trF_p^TF_e^TF_eSF_p^−TF_tr^T ·

³

b^(α)⊗d^(α)

´

+1 2

³

CA−(1 +δ_T)C^(α)´

E_e·(E_e−2B(θ−θ₀))

−CE_e·

³

B_A−(1 +δ_T)B^(α)

´

(θ−θ₀) ,

(18) 1

2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

For Peer Review Only

where B_A and B^(α) represent the thermal expansion tensors that include elastic stretch corrections of the austenite and the martensitic systemα, respectively, as given by

B_A= 1 2

¡(F_e^?)^TF_e^?A_A+A_A(F_e^?)^TF_e^?¢ , B^(α)= 1

2

³

(F_e^?)^TF_e^?A^(α)+A^(α)(F_e^?)^TF_e^?

´

. (19)

Note that (6) and (10)₂ have been used to get the expressions in (19). From (18) it can be observed that the bulk mechanical part of the transformation driving force consists of three terms, which successively represent the resolved stress, a term reflecting the difference in elasticity stiffness between the martensite product phase and the austenite parent phase, and a term reflecting their difference in thermal expansion properties. The dominant term in the mechanical driving force is typically the resolved stress whereas the two other terms play a role only when there is a relatively large difference in stiffness and/or thermal expansion between the parent and product phases. Furthermore, in accordance with [29], the thermal contribution to the transformation driving force, and the contributions related to the defect energy and surface energy, are derived as

f_th^(α)=ρ₀λ^(α)_T

θ_T (θ−θ_T) +ρ₀

³

h_A−h^(α)

´ µ

θ−θ_T−θln µ θ

θ_T

¶¶

, (20) f_d^(α)= ω_A

2

³

µ_A−(1 +δ_T)µ^(α)

´

β² and f_s^(α)= χ

`₀

³

2ξ^(α)−1

´

. (21) The dominant term in the thermal driving force is typically the thermal analogue of the resolved stress (i.e., the term related to latent heat) whereas the second term in (20) plays a role only when there is a relatively large difference in specific heats between the parent and product phases. Moreover, the contributions related to the defect energy and the surface energy play a role mainly at the nucleation stage of martensite (f_d effectively acting as a barrier if the martensite is stiffer than the austenite and f_s acting initially as a barrier to transformation due to the energy required to form new interfaces, see [37]).

The criterion for transformation nucleation and the evolution of the martensitic volume fraction in a transformation systemα are captured by the following kinetic relation:

ξ˙^(α)=







ξ˙₀tanh Ã

f^(α)−f_cr^(α) νf_cr^(α)

!

iff^(α)> f_cr^(α),

0 otherwise ,

(22)

wheref_cr^(α) is the critical value for the transformation driving force, ˙ξ₀ defines the maximum transformation rate andν is a viscosity-like parameter. The parameters ξ˙₀ and ν characterize the rate-dependent behavior of the above kinetic model.

For TRIP steels it is realistic to assume that plastic deformations evolve in the austenite only, and not in the martensite [35], i.e., the martensite deforms elastically. In the present study the evolution of plastic slip in the austenite is described by a power-law kinetic relation, where the rate of plastic slip in the 1

2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

For Peer Review Only

austenite, ˙γ_A⁽ⁱ⁾, depends on the corresponding driving force for plasticity,g⁽ⁱ⁾_A, as [29]

˙ γ_A⁽ⁱ⁾=









˙ γ₀^A



 Ãg⁽ⁱ⁾_A

s⁽ⁱ⁾_A

!_(1/pA)

−1



 ifg_A⁽ⁱ⁾> s⁽ⁱ⁾_A ,

0 otherwise ,

(23)

where ˙γ₀^A,p_Aand s⁽ⁱ⁾_A are the reference slip rate, the rate-sensitivity exponent and the resistance against plastic slip in slip systemi. Using the procedure presented in [29], the total driving force for plasticity in the austenite can be computed from the resolved stress contribution (and its thermal analogue) and the derivative of the Helmholtz energy with respect to the slip rate as

g⁽ⁱ⁾_A =F_e^TF_eS·

³

m⁽ⁱ⁾_A ⊗n⁽ⁱ⁾_A

´

+ρ₀θφ⁽ⁱ⁾_A −ω_Aµβw⁽ⁱ⁾. (24) Observe that the third term in the plastic driving force is always negative, hence it can also be interpreted as a hardening effect (i.e., an increase in difficulty to achieve plastic slip as more dislocations are generated).

The rate of change of the slip resistance, ˙s⁽ⁱ⁾_A, is constitutively connected to the rate of plastic slip in the austenite as [29]

˙ s⁽ⁱ⁾_A =

NA

X

j=1

H_A^(i,j)γ˙_A^(j). (25)

Here, H_A^(i,j) is the hardening matrix, which is characterized by self- and cross- hardening contributions, i.e.,

H_A^(i,j)= (

k^(j)_A fori=j,

q_Ak_A^(j) fori6=j, with k_A^(j)=k^A₀ Ã

1−s^(j)_A s^A_∞

!_uA

, (26)

whereq_A is the latent hardening ratio, which reflects the ratio between the cross- and self-hardening moduli on each slip system,k^A₀ is the reference hardening mod- ulus, s^A_∞ is the saturation value of the slip resistance, and u_A is the hardening exponent.

In equation (4) it has been assumed that the rate of change of the microstrain parameter, ˙β, is connected to the rate of change of plastic slip, ˙γ⁽ⁱ⁾, through the functionsw⁽ⁱ⁾. In principle, a purely kinematical model can be constructed for the functions w⁽ⁱ⁾. Alternatively, in the present study the kinetic model proposed by Clayton [36] is adopted, where the rate of change of the microstrain ˙β_A in the austenitic region is constitutively related to the average value of the rate of change of the slip resistance in the austenite, ˙s⁽ⁱ⁾_A, i.e.,

1 N_A

NA

X

i=1

˙

s⁽ⁱ⁾_A =c_Aµ_Aβ˙_A, (27)

withc_A a dimensionless scaling factor. From (4), (25), (27) and the relation ˙β = ξ_Aβ˙_A/J_tr, it follows that the functionsw⁽ⁱ⁾can be related to the hardening moduli 1

2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

For Peer Review Only

as

w⁽ⁱ⁾= 1 c_Aµ_AN_A

NA

X

j=1

H_A^(j,i). (28)

3. Thermo-elasto-plastic model for single-crystalline ferrite

The decomposition of the total deformation gradient and the entropy in the thermo- elasto-plastic model for single-crystalline ferrite can be formulated in a similar fashion as in the thermo-elasto-plasto-transformation model for austenite, given in Section 2, by suppressing the transformation contribution. Formally, this can be achieved by eliminating the volume fractions as state variables, and by setting F_tr=I andη_tr= 0. In addition, analogous to (5), the elastic deformation gradient F_eneeds to include the thermal expansion/contraction under a temperature change fromθ₀ toθ, in accordance with

F_e=F_e^?(I+A_F(θ−θ₀)) , (29) where F_e^? is the elastic deformation gradient measured at a temperature θ and A_F is the thermal expansion tensor, which is given by A_F = α_FI, with α_F the thermal expansion coefficient of ferrite. In analogy with equations (15) and (16), the mechanical and thermal constitutive relations, respectively, have the form

S =CF(E_e−B_F(θ−θ₀)) and η_e=h_Fln µ θ

θ_F

¶

+η_F+ 1

ρ₀CFE_e·B_F, (30) whereCF,ρ₀,h_F,θ_F andη_F are, respectively, the elastic stiffness tensor, the mass density, the specific heat, the reference temperature and the reference entropy of BCC ferrite. The stiffness tensor of BCC ferrite is characterized by three inde- pendent elastic moduli, κ^F_i(i= 1,2,3). In (30)₁, E_e represents the elastic Green- Lagrange strain andB_F is the thermal expansion tensor of ferrite with corrections for the elastic stretch, i.e.,

B_F= 1 2

¡(F_e^?)^TF_e^?A_F+A_F(F_e^?)^TF_e^?¢

. (31)

Furthermore, the asymmetry of slips in the twinning and anti-twinning directions, which is typical for plastic deformations in BCC crystals, can be accounted for by means of the “non-glide” stress effect, ˆτ_F⁽ⁱ⁾, defined by [38, 39]

ˆ

τ_F⁽ⁱ⁾=F_e^T F_eS·(m⁽ⁱ⁾_F ⊗nˆ⁽ⁱ⁾_F ) , (32) where ˆn⁽ⁱ⁾_F represents the non-glide plane corresponding to slip systemi. As pointed out in [29], the expression for the non-glide stress is formally similar to that of the resolved stress, with the normal to the non-glide plane playing an equivalent role as the normal to the slip plane. The non-glide stress is used for constructing the

“effective” slip resistance ˆs⁽ⁱ⁾_F as ˆ

s⁽ⁱ⁾_F =s⁽ⁱ⁾_F −ˆa⁽ⁱ⁾ˆτ_F⁽ⁱ⁾, (33) 1

2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

For Peer Review Only

Ferritic matrix (81 grains)

Austenitic inclusion (27 grains) f₁

f₂

f₃ Face 2 (u

1 = 0) Face 4

(u

3 = 0) Face 3

(u

2 = 0) Face 1 (u

1 = 10^-4at)

f₂ a f₃

Ferrite 〈100〉_F pole figure

f₂ f₃

Austenite 〈100〉_A pole figure

Figure 2. A polycrystalline TRIP steel sample and the applied boundary conditions that correspond to uniaxial tension. The orientations of the austenitic and ferritic grains are represented in terms ofh100iA- andh100iF-pole figures.

where s⁽ⁱ⁾_F is the common slip resistance, and ˆa⁽ⁱ⁾ is a weight parameter that de- termines the net contribution of the non-glide stress to the effective slip resistance of system i. The expressions for the driving force and kinetic relation for plastic slip in the BCC ferrite can be obtained analogous to those for the FCC austenite in (23) and (24), by applying the following replacements

g_F⁽ⁱ⁾→g⁽ⁱ⁾_A , m⁽ⁱ⁾_F →m⁽ⁱ⁾_A , n⁽ⁱ⁾_F →n⁽ⁱ⁾_A , ˆs⁽ⁱ⁾_F →s⁽ⁱ⁾_A , γ˙₀^F →γ˙₀^A, p_F →p_A, φ⁽ⁱ⁾_F →φ⁽ⁱ⁾_A , ω_F →ω_A, µ_F→µ, β_F →β, and w⁽ⁱ⁾_F →w⁽ⁱ⁾.

Finally, the evolutions of the slip resistances⁽ⁱ⁾_F and the microstrainβ_Fare described by similar laws as those used for the FCC austenite, see equations (25)–(28).

4. Simulation of TRIP-assisted steel behavior under thermomechanical loading

4.1. Microstructural sample and material parameters

In this section, simulations are performed on microstructural TRIP steel samples subjected to thermomechanical loading. The microstructural samples contain 27 polyhedral grains of (retained) austenite with random crystallographic orienta- tions. All austenitic grains have the same volume and the same carbon concen- tration (i.e., 1.4 wt. %), and are distributed uniformly in a matrix constructed of 81 randomly-oriented grains of ferrite, see Figure 2, where the random crystallo- graphic orientations assigned to the austenitic and ferritic grains are represented in h100i_A and h100i_F pole figures, respectively. The grains of retained austenite occupy 13 % of the initial, undeformed sample volume.

Observe that the sample is only qualitatively representative of a TRIP mi- crostructure since, in an actual sample, the grains are typically not randomly oriented and some (spatial) variations are observed in grain sizes and carbon concentrations. In order to account for microstructural texture while preserving a reasonable computational time, an additional homogenization step is required (see, e.g., [40] for TRIP steels in a purely mechanical context). The simulations presented here, however, focus on interactions at the grain level.

The single-crystalline thermo-elasto-plastic-transformation model described in 1

2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

For Peer Review Only

Table 1. Material parameters for the austenite/martensite model and the ferrite model.

Austenite/martensite Values

Elastic moduli austenite κ^A₁ = 286.8,κ^A₂ = 166.4,κ^A₃ = 145.0 [GPa]

Elastic moduli martensite κ^M₁ = 372.4, κ^M₂ = 345.0,κ^M₃ = 191.0, κ^M₄ = 508.4, κ^M₅ = 201.9,κ^M₆ = 229.5 [GPa]

Transformation shape strain and dilation γ_T =kb^(α)k = 0.1809, δ_T=b^(α)·d^(α) = 0.0391 Transformation kinetic law ξ˙₀= 0.003 s⁻¹,ν = 0.17,f_cr^(α)= 227 MPa Surface energy χ= 0.2 J·m⁻²,`₀ = 0.05 µm

Hardening law s_A,0 = 189, s^A_∞= 579 [MPa], k₀^A= 3 GPa,u_A= 2.8,q_A= 1 Plasticity kinetic law γ˙₀^A= 0.001 s⁻¹,n_A= 0.02 Defect energy β_A,0 = 0.0056,c_A= 0.5,ω_A= 10

µ_A= 67.5,µ^(α)=µ_M= 98.4 [GPa]

Thermal driving forces λ^(α)_T =−50.5 kJ·kg⁻¹,φ⁽ⁱ⁾_A = 5.13 m²·K⁻¹·s⁻² Coefficient of thermal expansion α_A=α_M= 2.1×10⁻⁵ K⁻¹

Ferrite Values

Elastic moduli ferrite κ^F₁ = 233.5,κ^F₂ = 135.5,κ^F₃ = 118.0 [GPa]

Thermal driving force φ⁽ⁱ⁾_F = 4.27 m²·K⁻¹·s⁻² Hardening law s_F,0 = 154, s^F_∞= 412 [MPa]

k₀^F= 1.9 GPa, u_F = 2.8,q_F = 1 Plasticity kinetic law γ˙₀^F = 0.001 s⁻¹,n_F= 0.02 Non-glide stress coefficient ˆa⁽ⁱ⁾= 0.12

Defect energy ω_F= 7, β_F,0 = 0.0056,c_F = 0.5,µ_F = 55.0 GPa Coefficient of thermal expansion α_F = 1.7×10⁻⁵ K⁻¹

Section 2 is used to simulate the response of the austenitic region whereas the thermo-elasto-plasticity model presented in Section 3 is applied to the ferritic re- gion. The material parameters of these models are listed in Table 1. The values in this table are equal to those used in [29], except for the critical value of the trans- formation driving force fcr^(α) and the latent heat λ^(α)_T , which are taken somewhat different. The differences stem from a reinterpretation of these parameters in order to realistically simulate the onset of transformation in coupled thermomechanical problems and to adequately account for the thermomechanical coupling during transformation. Furthermore, the elasto-plastic properties of the material, e.g., the elastic moduli, the initial slip resistance and the hardening moduli, are assumed to be temperature-independent for simplicity reasons. Details on the computation of the vectors b^(α) and m^(α) of the 24 martensitic transformation system and a complete list of these vectors can be found in previous works [26–28]. The slip directions m⁽ⁱ⁾_F , the slip plane normals n⁽ⁱ⁾_F , and the non-glide plane normals ˆn⁽ⁱ⁾_F of the 24 slip systems in the BCC ferrite model are listed in [30].

The entropy density change per unit slip in the austenite,φ⁽ⁱ⁾_A, and in the ferrite, φ⁽ⁱ⁾_F , are assumed to be temperature-independent. The values of these parameters, as presented in Table 1, are consistent with the values of the thermal driving force for plasticity presented in [29], i.e.,φ⁽ⁱ⁾is computed from the value ofρ₀θφ⁽ⁱ⁾ given in [29], evaluated at θ = θ₀ = 300 K where the mass density in the reference configuration for the austenite and ferrite isρ₀= 7800 kg·m⁻³. The specific heats of the austenite and the martensitic transformation systems are assumed to be similar, i.e.,h_A≈h^(α) for allα = 1, . . . , M. Correspondingly, the thermal part of 1

2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

For Peer Review Only

the transformation driving force in the austenite, as represented by (20), can be approximated as

f_th^(α)(θ)≈ρ₀λ^(α)_T

θ_T (θ−θ_T) . (34)

The theoretical transformation temperature is estimated asθ_T ≈633 K [27]. The value of the latent heat λ^(α)_T for the transformation from austenite to martensite (related to the temperature θ_T) is assumed to be equal for all transformation systemsα, i.e.,λ^(α)_T =λ_T. Furthermore, for simplicity, the values of the coefficients of thermal expansion of the austenite and martensite are assumed to be the same, α_A = α_M, and, together with the value for ferrite, α_F, are within the range of values reported in the literature (see, e.g., [41]).

4.2. Simulation of TRIP steel sample undergoing uniaxial tensile loading at various temperature

The sample is initially stress-free and is subsequently subjected to an isothermal uniaxial tensile loading during a time interval 0< t≤1000 s, as prescribed through the following boundary conditions, see also Figure 2: (i) The normal displacement and the tangential tractions on the external faces 2, 3 and 4 are set equal to zero;

(ii) The normal displacement on the external face 1 is prescribed asu₁ = 10⁻⁴at, where a is the side-length of the cubic sample, and the tangential tractions are set to zero; (iii) the remaining cube faces (5 and 6) are assumed to be traction- free. The applied boundary conditions correspond to a maximum axial nominal strain of 0.1, which develops at a constant straining rate of 10⁻⁴ s⁻¹. In order to study the influence of temperature on the overall mechanical response of TRIP steels, the isothermal uniaxial tensile simulations are performed at three different temperatures, i.e., 350, 300 and 250 K.

Figure 3 depicts, respectively, the evolution of the austenitic volume fraction as a function of the average axial strain (Figure 3a), the average axial stress- strain response (Figure 3b) and the axial stress-strain responses of the individual austenitic/martensitic (Figure 3c) and ferritic (Figure 3d) phases. The effective responses in the latter two figures are obtained by averaging the local stresses and strains in the grains of austenite (which may contain martensite) and ferrite, respectively.

The numerical simulations are performed with the finite element pro- gram ABAQUS, where the martensitic transformation model for the austenitic/martensitic phases and the crystal plasticity model for the ferritic phase are implemented as a so-called “user subroutine”. The numerical implementation of the transformation model is based on the formulation presented in [28], where the incremental update algorithm relies on a fully implicit Euler backward dis- cretization and a consistent tangent operator. A robust search algorithm is used for detecting the transformation systems that are activated during thermomechan- ical loading. Although not described in [28], a similar search algorithm is used to identify the active slip systems in the austenitic phase. In addition, the nu- merical implementation of the crystal plasticity model for the ferritic phase is based on the formulation proposed in [42], where improvements taken from the numerical implementation of the transformation model [28] have been incorpo- rated. The sample shown in Figure 2 has been discretized with approximately 8000 tetrahedral elements (about 7000 for the 81 ferritic grains and 1000 for the 27 austenitic/martensitic grains). Each grain has been assigned one of the crystal 1

2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

For Peer Review Only

150 0 900 750 600 450 300

0 0.04 0.06 0.08 0.10

0 -200 1000 800 600 400 200

(1) θ = 350 K (2) θ = 300 K (3) θ = 250 K

(1)

(d) (c)

Cauchy stressT11 ofaust./mart. [MPa]A

Logarithmic strain e₁₁ of ferrite 0 0.04 0.06 0.08 0.10 0.02

Logarithmic strain e₁₁ of austenite 0.02

(2) (3)

(1) (3) (2)

(1) θ = 350 K (2) θ = 300 K (3) θ = 250 K Cauchy stressT11 of ferrite [MPa]F

A F

0.025 0 0.150 0.125 0.100 0.075 0.050

0 0.04 0.06 0.08 0.10

Austenitic volume fractionξA_

Logarithmic strain e₁₁_ 0.02

(1) (2)

(3) (1) θ = 350 K

(2) θ = 300 K (3) θ = 250 K

150 0 900 750 600 450 300

(1) θ = 350 K (2) θ = 300 K (3) θ = 250 K

(1)

Cauchy stressT11 [MPa]_

0 0.04 0.06 0.08 0.10

Logarithmic strain e₁₁_ 0.02

(2) (3)

Figure 3. TRIP steel samples undergoing uniaxial tensile loading at various temperatures: (a) Evolution of austenitic volume fraction as a function of axial strain, (b) effective axial stress-strain response of the sample, (c) effective axial stress-strain response of the austenitic/martenstic phases, and (d) effective axial stress-strain response of the ferritic phase.

orientations shown in the pole figures in Figure 2

As shown in Figure 3a, under mechanical loading the transformation rate of austenite depends strongly on the temperature at which the transformation takes place. At a relatively low temperature of 250 K the transformation is virtually completed at about 0.04 axial deformation. In contrast, at a relatively high tem- perature of 350 K only a small portion of retained austenite has transformed into martensite at the end of the loading process (i.e., at 0.1 axial deformation). For an intermediate temperature of 300 K, about half of the austenite has transformed into martensite at the maximum imposed strain. The dependency of the transformation rate on the temperature is due to the following two factors: Firstly, in view of (34) and sinceλ_T is negative, the thermal transformation driving force, f_th^(α) decreases with temperature, (i.e., austenitic grains are more stable at a higher temperature), which thus makes the transformation easier at lower temperature. Secondly, the thermal contribution to the plasticity driving force,ρ₀θφ⁽ⁱ⁾_A, increases with temper- ature, meaning that plastic flow becomes more dominant at elevated temperature.

Since the development of plastic deformations retards the transformation process, the transformation occurs faster at lower temperature. As a consequence of these two factors, in the simulation at low temperature, θ = 250 K, the dominant in- elastic mechanism in the austenite is the martensitic transformation, whereas the dominant mechanism in the simulation at high temperature,θ= 350 K, is plastic flow.

When comparing cases 1 and 3 in Figure 3b, it can be observed that a relatively 1

2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60