Temperature-dependent micromechanical behavior of medium-Mn transformation-induced-plasticity steel studied by in situ synchrotron X-ray diffraction

Accepted Manuscript Temperature-dependent micromechanical behavior of medium-Mn transformationinduced-plasticity steel studied by in situ synchrotron X-ray diffraction Minghe Zhang, Longfei Li, Jie Ding, Qingbao Wu, Yan-Dong Wang, Jonathan Almer, Fangmin Guo, Yang Ren PII:

S1359-6454(17)30764-4

DOI:

10.1016/j.actamat.2017.09.030

Reference:

AM 14060

To appear in:

Acta Materialia

Received Date: 10 June 2017 Revised Date:

30 August 2017

Accepted Date: 14 September 2017

Please cite this article as: M. Zhang, L. Li, J. Ding, Q. Wu, Y.-D. Wang, J. Almer, F. Guo, Y. Ren, Temperature-dependent micromechanical behavior of medium-Mn transformation-inducedplasticity steel studied by in situ synchrotron X-ray diffraction, Acta Materialia (2017), doi: 10.1016/ j.actamat.2017.09.030. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

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Graphical abstract

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ACCEPTED MANUSCRIPT Temperature-dependent

micromechanical

behavior

of

medium-Mn

transformation-induced-plasticity steel studied by in situ synchrotron X-ray diffraction

Almerc,*, Fangmin Guoc, Yang Renc

State Key Laboratory for Advanced Metals and Materials, University of Science and

Technology Beijing, Beijing 100083, China

Collaborative Innovation Center of Steel Technology, University of Science and

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b

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a

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Minghe Zhanga,b,c, Longfei Lia, Jie Dinga, Qingbao Wua, Yan-Dong Wanga, , Jonathan

Technology Beijing, Beijing 100083, China c

X-ray Science Division, Advanced Photon Source, Argonne National Laboratory,

Argonne, IL 60439, USA

Abstract

temperature-dependent

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The

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*Corresponding author Tel: +8601082377942; E-mail: [email protected] (Yan-Dong Wang) or [email protected] (Jonathan Almer)

micromechanical

behavior

of

medium-Mn

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transformation-induced-plasticity (TRIP) steel with a nominal chemical composition of Fe-0.1C-10Mn-2Al (mass%) fabricated by intercritical annealing 650

for 1h

after cold-rolling, was investigated using in situ high-energy X-ray diffraction (HE-XRD) with uniaxial tensile tests at temperatures of 100, 25 and -50

. We find

that as the deformation temperature decreases, the Lüders strain decreases and more austenite transforms to martensite during Lüders band propagation. The Lüders bands are associated with sudden changes of lattice strain in austenite. At 100 1

, austenite is

ACCEPTED MANUSCRIPT too stable to transform to martensite, resulting in limited work-hardening capability and a relatively low strain to failure. At 25

, the austenite is found to transform in

bursts during applied loading. These transformations correlate with stepwise peak

band propagation. At -50

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broadening in the austenite phase and are attributed to Portevin-Le Châtelier (PLC) , we observe a more intense TRIP effect which suppresses

PLC band formation and leads to a high ultimate tensile strength. In addition to the

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TRIP effect, we find that work hardening in the deformation-induced martensite phase

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plays an important role in the plastic stability at low temperature, by accommodating stress compatibility between grains and phases. Our studies provide the first direct experimental evidence on the existence of large stress concentrations in the austenite phase near the Lüders band propagation front. These observations help to deepen

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understanding of the complex temperature-dependent micromechanical behaviors of advanced medium-Mn TRIP steels.

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Keywords: Medium-Mn TRIP steel; High-energy X-ray diffraction; Austenite stability;

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Load partitioning

1. Introduction

Advanced high strength steels (AHSS) that demonstrate an excellent combination of high strength and plasticity are the subject of intense research and development in recent years, owing to an increasing demand for fuel efficiency, emissions reduction and safe design in the automobile industry [1-3]. In comparison with the first 2

ACCEPTED MANUSCRIPT generation AHSS, such as dual phase (DP) steel and transformation induced plasticity (TRIP) steel, and the second generation AHSS, such as twinning induced plasticity (TWIP) steel, the third generation AHSS provides a good trade-off between the

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mechanical properties and the cost of materials fabrication and final production [4-7]. Recently, medium-Mn TRIP steels have been developed [8-11], which contain 5 to 10 mass% Mn with fine-grained duplex ferrite-austenite microstructure and demonstrate

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a combination of high ultimate tensile stress (UTS) and total elongation (TE) due to

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the TRIP effect. Previous studies have revealed the effects of alloying elements and intercritical annealing conditions on the microstructure and mechanical properties of such steels [12-15]. For Fe-0.11C-5.7Mn (mass%) steel with mixed microstructure of ferrite and 29 vol.% retained austenite (RA), its UTS of 878 MPa and TE of 34% can

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be reached as reported by Miller in 1972 [16]. Luo et al. [17] reported that Fe-0.2C-5Mn (mass%) steel exhibited UTS of 785–1100 MPa with a TE of 20–30%. Yen et al. [18] achieved a yield strength (YS) of 1080 MPa, UTS of 1390 MPa and TE

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of 26.7% with Fe-10.6Mn-0.08C, which was obtained by applying a short annealing

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at 610 °C for 8 min. Excellent mechanical properties were obtained with higher C and Mn concentrations, which is attributed to an increase in the volume fraction of RA in the medium-Mn TRIP steels. Furthermore, the addition of Al in medium-Mn TRIP steels can suppress the formation of cementite and increase the intercritical annealing temperature [19]. Cai et al. [20] reported that a 0.2C-11Mn-4Al steel with about 67 vol.% RA, obtained by annealing at 770 °C for 3min, demonstrated the tensile strength and uniform elongation up to about 1007MPa and 65%, respectively. A 3

ACCEPTED MANUSCRIPT Fe-0.2C-8Mn-5Al steel exhibited a high tensile strength of 949 MPa and a TE of 54% after annealing at 900 °C for 2min [21]. However, the load partitioning among constituent phases in medium-Mn TRIP steels as a function of applied

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thermos-mechanical load is still not fully understood. Such information promises to further improve and/or tailor the mechanical properties of these steels.

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Neutrons and high-energy X-ray diffraction (HE-XRD) each provide bulk probes

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which can be used for in situ thermo-mechanical characterization of a variety of materials [9, 22-26]. Muránsky et al. [23] investigated the deformation behavior of low-alloyed TRIP-assisted steels by in-situ neutron diffraction during tensile tests at room temperature and two elevated (50

and 100

) temperatures. They found that

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the applied tensile load is redistributed within the complex TRIP-steel microstructure and the deformation-induced martensitic transformation of retained austenite dispersed within the matrix governs the plastic deformation of TRIP-assisted steels.

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Blonde et al. [25] studied the stability of the RA in low-alloyed TRIP steels using in . It was

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situ HE-XRD during tensile tests at variable temperatures down to -120

shown that at lower temperatures the deformation-induced austenite transformation is significantly enhanced, resulting in a higher elongation at fracture, and the mechanical stability of RA in TRIP steel is found to be a complex interplay among carbon concentration in the austenite, grain orientation, load partitioning and temperature. Fu et al. [26] conducted research on the stress partitions among multiple phases for two low-alloyed TRIP-assisted steels, using in situ HE-XRD at ambient and low 4

ACCEPTED MANUSCRIPT temperatures (-40

) under uniaxial tensile loading. Significant stress partitioning

was observed among phases during plastic deformation and retained austenite undertook the largest load, followed by bainitic ferrite, while the ferrite matrix phase

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bore the smallest load. To date, however, the micromechanical response of medium-Mn TRIP steels at variable temperatures and applied loads has not been

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transformation and resulting load partitioning.

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systematically evaluated, particularly in the relation to the interplay between phase

In this study, in situ HE-XRD experiments were performed to study the micromechanical behaviors of a medium-Mn TRIP steel during tensile loading at 100, 25 and -50

, respectively, to elucidate the temperature-dependent deformation

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characterizations. The kinetics of deformation-induced transformation of RA to martensite, evolution of lattice strain and peak broadening of constituent phases are presented. The load partitioning among constituent phase and the interplay among

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Lüders band, Portevin-Le Châtelier (PLC) band, and TRIP effect is discussed in detail.

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This study provides a better understanding of deformation mechanisms of medium-Mn TRIP steels at various deformation temperatures and promises to enhance the design and optimization of medium-Mn TRIP steels.

2. Materials and experimental procedures The Fe-0.1C-10Mn-2Al steel was fabricated by vacuum induction melting, and it’s resulting chemical composition is presented in Table 1. After being homogenized at 5

ACCEPTED MANUSCRIPT 1200 °C for 1 h, 20 mm thick plates were hot-rolled between 1100 and 900 °C to 4 mm thickness and then cold-rolled to 1.5 mm thickness. Specimens for microstructural observations and tensile tests were cut from thin plates by

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electrosparking and subsequently annealed at 650 °C for 1 h and then air-cooled. Microstructural observation was performed using a ZEISS SUPRA55 field-emission scanning electronic microscope (FE-SEM) after etching with 4 vol.% Nital. The Mn

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and Al contents of the austenite phase were determined using energy-dispersive

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spectroscopy (EDS) in the scanning electron microscope. The average sizes of the ferrite grains and retained austenite grains were measured using Image-Pro Plus 6.0 (produced by Media Cybernetics Company, USA) image-analysis software by means of linear intercepts taken in the SEM micrographs. The carbon content and volume

and [28], respectively. In

situ

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fraction of retained austenite were determined by the methods depicted in Ref. [27]

synchrotron-based

HE-XRD

technique

was

used

to

study

the

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temperature-dependent deformation behavior of experimental steel on the 11-ID-C

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beam line of the Advanced Photon Source, Argonne National Laboratory, USA. A monochromatic X-ray beam with an energy of ∼105 keV (with wavelength 0.11742 Å) was used to study the microstructure evolution and change in lattice strain for specimens under the tensile loading mode, with the experimental set-up shown schematically in Fig. 1. The dog-bone-shaped tensile specimen with the gauge part of 10 mm (length) × 3 mm (width) × 0.5 mm (thickness) was mounted on a stress rig with the rolling direction parallel to the loading direction (LD). Two-dimensional 6

ACCEPTED MANUSCRIPT (2-D) diffraction patterns, composed of Debye rings, under different loading levels were collected during tensile straining at selected temperatures of 100, 25 and -50

.

The data from the section with the scattering vector parallel to the LD were integrated

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for calculating the tensile lattice strains for different phases. The tensile lattice strain ε of plane (hkl) with the normal paralleling the LD can be determined by the relative change in the corresponding lattice plane spacing, Q (hkl ) − Q0 (hkl ) Q0 (hkl )

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ε=−

(1)

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where Q0 (hkl) and Q (hkl) are the modulus of scattering vector of hkl-crystal plane before and after an applied stress σ at the test temperature. The full width at half-maximum (FWHM) for different diffraction peaks was used for quantifying strain heterogeneity, which was obtained through correcting the instrumental

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broadening with a near-perfect (broadening-free) CeO2 sample (Gaussian peak shape

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is assumed here) [29, 30].

3. Experimental results

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3.1 Microstructures and mechanical properties of Fe-0.1C-10Mn-2Al steel The microstructure of the studied steel fabricated by annealing at 650

for 1 h after

cold rolling consists of ultrafine grained austenite (γ) and ferrite (α), as shown in Fig. 2. The intercritical annealing induces an obvious reduction in grain size for both austenite and ferrite, exhibiting the average size of about 0.73µm and 0.47µm, respectively. The average values of Mn and Al contents for austenite are 13.34 mass% and 1.55 mass%, respectively, which were obtained through the SEM/EDS 7

ACCEPTED MANUSCRIPT measurements of more than 25 randomly-selected austenite grains. The estimated carbon concentration XC of austenite is 0.18 mass% by the methods depicted in Ref.

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[27].

The in situ tensile stress vs. strain curves of the investigated steels are shown in Fig. 3, with the mechanical properties detailed in Table 2. All curves existed an upper yield

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point (UYP) and a lower yield point (LYP), followed by Lüders straining. With , YS of the studied steel

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decreasing the deformation temperature from 100 to -50

showed a slight increase from 940 MPa to 976 MPa, similar to the expected temperature-dependent strength of steels [31]. After entering into the stage of Lüders straining, with decreasing the deformation temperature from 100 to -50

, the

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mechanical behavior changed dramatically, characterized by a remarkable change in the UTS and TE, as shown in Fig.3. Such temperature-dependent plastic behavior found in the studied steel was not only completely different from that expected for

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metals and multiphase alloys with dislocation glide and dislocation interactions as the

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main mechanism to control the plastic flow and work-hardening [31], but also somewhat different from that for low-alloyed TRIP steels deformed at temperatures either above or below room temperature [23, 25]. The reason for the different temperature-dependent mechanical behavior for the Fe-0.1C-10Mn-2Al steel and the low-alloyed TRIP steel may be attributed to that the higher Mn content causes an enhancement in temperature-dependent mechanical stability of austenite.

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ACCEPTED MANUSCRIPT 3.2 Evolution of austenite volume fraction during deformation at different temperatures The 2D HE-XRD diffraction image covering over the entire 360° azimuthal range

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recorded for the studied steels deformed at three selected temperatures are shown in , all diffraction peaks for γ phase were

Fig.4. For the sample deformed at 100

changed slightly even near fracture. While, for the sample deformed at 25 or -50

,

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all diffraction peaks for γ phase became weak gradually with increase plastic strain

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and even some peaks finally disappeared at large plastic strains.

Fig.5 shows the change in volume fraction of RA in the studied steel during tensile deformation at above-mentioned three temperatures. Before loading, the initial

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volume fraction of RA in the used steel is 76%, 76% and 74% for the sample deformed at 100, 25 and -50

, respectively, derived from the diffraction patterns.

The minor difference in the volume fraction of the initial RA at the different

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temperatures was caused by a possible tiny transformation from austenite

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(γ)-to-martensite (α') under cooling. For the sample deformed at 100

, only a small

part of RA disappeared until fracture. For the sample deformed at 25 or -50

, the

volume fraction of RA decreased suddenly at a strain of about 0.05 after sample yielding, which was in the half of Lüders straining as shown in Fig.3. This phenomenon indicates that the typical strain-induced martensite were generated during deformation. With increasing plastic strain, the volume fraction of RA decreased step by step in the sample deformed at 25 9

and decreased gradually in that

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. The difference in change of volume fraction of RA during tensile loading at

different temperatures, owing to the temperature-dependent stability of austenite, resulted in the strongly temperature-dependent mechanical behaviors of the used steel

3.3 Evolution of lattice strain at different temperatures

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as shown in Fig.3.

Synchrotron-based HE-XRD having quite high reciprocal space resolution makes it

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possible to separate the overlapped peaks of ferrite and martensite, despite their

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similar crystal structures. The α-211 and γ-311diffraction peaks were chosen for calculating the body-centered cubic (bcc) and face-centered cubic (fcc) lattices, respectively, due to their insensitivity to intergranular strain [32, 33].

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Fig.6 shows the evolution of the α-211 and γ-311diffraction patterns during deformation at 100, 25 and -50

, respectively. It can be seen clearly that after the

samples yielding at the above-mentioned temperatures, with increasing the applied

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loading, both diffraction intensity of γ-311peak or α-211 peak decreased and their

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diffraction peak positions are shifted toward the low Q value. For the sample deformed at 100

, the α-211 peak shape changed slightly after yielding. In contrast,

for the sample deformed at 25 or -50

, the α-211 peak shape became asymmetry

after yielding, due to the formation of α'-martensite that has similar lattice with ferrite. According to the method proposed by Jia et al. [22] and Yan et al. [24], the α-211 diffraction peaks with the scattering vector parallel to the LD were fit with

two

Gaussian functions characterized as the ferrite and martensite, respectively, as 10

ACCEPTED MANUSCRIPT illuminated in Fig.7a and Fig.7.b. With increasing the applied stress, more RA transformed to martensite, leading to an increase in peak intensity for the martensite. Moreover, due to the initial martensite transformed from austenite may bear some

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load during tensile deformation, d0 of α'-211 as reference state could be obtained by plotting the d-spacing vs. sin2 Ψ, as shown in Fig.7c, where Ψ is the azimuth angle

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relative to the loading direction by cos Ψ = cos η cos θ [24, 34].

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The lattice strains εhkl along LD for different {hkl} planes of the constitute phases of the studied steel are plotted as a function of the applied stress and strain in Fig. 8. For the sample deformed at 100

, the lattice strain of ferrite or austenite increased firstly

during elastic straining and decreased slightly after yielding, showing a similar

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tendency of variation. Moreover, the difference of lattice strain of austenite and ferrite indicated that load partitioning took place after yielding. At a larger Lüders strain, the lattice strain of ferrite increased obviously while the lattice strain of austenite changed

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jaggedly. With further increasing deformation strain, the lattice strain of ferrite or

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austenite increased gradually with straining until necking. Overall, the sample has a relative low work-hardening rate, resulting in the ultimate tensile strength smaller than the yielding strength as shown in Fig. 3.

For the samples deformed at 25 and -50 , load partitioning between austenite and ferrite also took place after yielding. In the beginning of Lüders straining, the evolution of the lattice strain of ferrite or austenite looks similar to that deformed at 11

ACCEPTED MANUSCRIPT 100

. However, due to the formation of martensite at a larger Lüders strain, the

lattice strain of ferrite or austenite decreased to some extent. Subsequently, after the beginning of the plastic deformation, the lattice strains of austenite and martensite

Particularly,

the

deformation-induced

work-hardening ability at 25 and -50

martensite

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increased markedly while the lattice strain of ferrite increased gradually until necking. exhibits

the

remarkable

, which could provide the stress gradient to

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accommodate the stress/strain compatibility between grains and phases (f.g, ferrite

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and austenite). We thus believe that, besides effect of the transformation kinetics, the remarkable localized work-hardening ability due to formation of deformation-induced martensite contributes to the plastic stability at low temperature in the medium-Mn TRIP steel. It should be noted that for the sample deformed at 25

, the zigzag

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change in the lattice strain of ferrite was corresponding to a kind of discontinuous transformation of martensite. It is very interesting to find that there is a sudden change

Section 4.

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in lattice strain of γ-311 during deformation, which will be explained in detail in the

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3.4 Peak broadening of the constitute phases at different temperatures Besides the change in the diffraction intensity and peak position, the peak broadening for austenite, ferrite and martensite during plastic deformation provides further information on evolution of microstructures, which is shown in Fig.6. Such broadening of diffraction peaks can be attributed to the inhomogeneous microstrain induced by dislocation-like defects (i.e. strain broadening) or the shrinkage of the coherent scattering volume (i.e. size broadening), or both of them [29]. The full width 12

ACCEPTED MANUSCRIPT at half maximum (FWHM) (denoted there as ∆Q/Q) of diffraction peak is usually used to characterize the extent of peak broadening. The FWHM of physical line-shape for the studied samples was calibrated by extracting the instrumental broadening from

[

]

1 2 2

FWHM = (FWHM measured ) − (FWHM instrumental ) 2

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the sample (measured) broadening via the following equation, i.e. (2)

Fig.9 shows the change of FWHM of constituent phases as a function of applied strain

For the sample deformed at 100

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, respectively.

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for the samples deformed at 100, 25 and -50

, the FWHM of α-211 or γ-311increased slowly in

the elastic straining stage and changed slightly in the Lüders straining stage. Subsequently, the FWHM of α-211 or γ-311increased markedly corresponding to the

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remarkable change in the lattice strain shown in Fig.8b. Moreover, the FWHM of α-211 or γ-311increased gradually with increasing deformation until necking. Overall, the FWHM value of γ-311was higher than that of α-211. For the samples deformed at , the variation of FWHM as a function of applied strain for α-211 or

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25 and -50

, but the FWHM values of γ-311 were

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γ-311was similar to that deformed at 100

much higher than that of α-211. The low temperature deformation indeed induces an obvious increase in FWHM for γ-311. It can be also evidenced that there is a stepwise change in FWHM for γ-311 for the sample deformed at 25

, which coincides exactly

with the abrupt change of volume fraction of RA. For the samples deformed at 25 and -50

, there is a similar value in FWHM for α'-211.

4. Discussion 13

ACCEPTED MANUSCRIPT For medium-Mn TRIP steels, the majority of the research efforts have so far focused on the overall influence of the volume fraction of metastable austenite as well as its stability on the strength and ductility. In the present study, metastable austenite also

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played an important role in the temperature-dependent mechanical behaviors of the used steel, as mentioned above.

,–the stability of metastable RA was very high and

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For the sample deformed at 100

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the yield stress is not high enough to trigger martensitic transformation, thus the mechanical behavior of the investigated steel was very similar to that of ultrafine-grained steels [10]. Due to the lack of work-hardening capability, the tensile straining of ultrafine grained steels is commonly considered to realize through the

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propagation of Lüders bands, as confirmed by Lee et al in ultrafine-grained Mn TRIP steel using in situ infrared (IR) thermography [10]. On the other hand, in the present work, the evolution of the FWHM of α-211 or γ-311, as shown in Fig. 9a, indicated

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that the accumulation of dislocations in austenite was enough to allow the used steel

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to exhibit a relative low work-hardening capability at the end part of Lüders straining. Thus, the sample deformed at 100 ℃ demonstrated still some extent of plastic deformation, as shown in Fig.3.

For the sample deformed at 25 ℃,–the stability of metastable RA was decreased with the temperature and deformation-induced transformation from RA to martensite could be triggered under a proper external stress or strain. The first decrease in the volume 14

ACCEPTED MANUSCRIPT fraction of RA in the early stage of Lüders straining, i.e. the occurrence of deformation-induced martensitic transformation, should be attributed to the local deformation caused by the propagation of Lüders band. Such Lüders band-promoted transformation

phenomenon

was

also

observed

in

metastable

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martensitic

nanostructured austenitic steel [34] and other medium Mn TRIP steels [35, 36] at room temperature. In order to illuminate the sudden change of lattice strain and the

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stepwise transformation kinetics of RA, we performed the X-ray scanning in the

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sample gauge section along LD with a total 18 scan points in each displacement step size of 300 µm. The change of volume fraction, FWHM and lattice strain of austenite at different applied strain levels over the effective sample gauge can be traced by the above-mentioned X-ray line-scanning along LD, which is shown in Fig.10. Once

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Lüders band propagates at some point, austenite transforms in bursts to martensite, associating average (cell) volume expansion of from γ-to-α'. Besides this, with the formation of martensite, the load partitioning by transformed martensite increases.

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Such cell volume expansion and load partitioning certainly lead to the sudden

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decreases of lattice strain of RA. After Lüders band propagating over the entire gauge, the stepwise transformation kinetics of austenite correlated exactly with stepwise peak broadening in the austenite, which can be observed in Fig.5, Fig.9b and Fig.11. This phenomenon could be explained by another plastic instability, namely Portevin-Le Châtelier (PLC) band, which already have been observed in other medium-Mn TRIP steels [37, 38]. As these bands nucleated at the specimen end and exhibited a continuous propagating feature, displayed in Fig.5 and Fig.11, they should be 15

ACCEPTED MANUSCRIPT characterized as the Type-A PLC band [39]. For the sample deformed at 25

, PLC

band propagation may promote about 10% volume fraction of phase transformation from RA to martensite. The detailed analysis from the in-situ HE-XRD data for

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revealing the real physical mechanisms on formation of the Type-A PLC band in the investigated steels is still in progress now.

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Moreover, further detailed analysis on the lattice strain distributions over the whole

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gauge section (as shown in Fig. 10c) at the applied strain of 0.07 and 0.1 reveals that there is a large strain gradient in austenite near the front of Lüders band propagation. The measured large lattice strain indicates the existence of highly-localized stress field in the austenite. The conventional theory propagative Lüders band focused

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mainly on localization temporal evolution of microstructures without any spatial dependence. The detailed scenario on the Lüders propagation related to the materials-length-scale gradient in stress/strain field is so far still missing, in spite of

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great efforts toward understanding the Lüders process from the thermal aspect [37].

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Here we give the direct experimental evidence on the existence of high stress concentration near the front of Lüders band, which is likely due to the formation of stable C-rich austenite associated with the redistribution of carbon after the partial martensitic transformation. Usually a very low or zero strain hardening may lead to specimen failure during plastic deformation. However, the studied medium-Mn TRIP steels undergo stable plastic deformation, which is still mysterious to materials science. The high stress concentration observed in the austenite certainly provides an 16

ACCEPTED MANUSCRIPT important mechanism on the stable propagation of Lüders band, i.e. the localized work-hardening capability on the front of Lüders band.

, more volume fraction of RA transforms to

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For the sample deformed at -50

martensite during Lüders band propagating, leading to the larger decreases in lattice strain of austenite. The stability of metastable austenite in the sample deformed at , was lower than that in the sample deformed at 25

, resulting in the gradual

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-50

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transformation of austenite in the former condition rather than the stepwise transformation of austenite at higher temperature during the stage of uniform plastic deformation. This indicates no PLC band propagates at -50

. This phenomenon is

explained here in terms of two factors: (1) With the formation of more volume

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fraction of martensite during Lüders band propagation associating more volume expansion of transformation, more dislocations accumulate in austenite that was sufficient to generate enough work-hardening capability in the final stage of Lüders

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straining. (2) The deformation-induced martensitic transformation is significantly

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enhanced at lower temperature, which might have a suppression effect on PLC band. 5. Conclusions

Ultrafine-grained austenite and ferrite microstructure with the volume fraction of austenite of about 76% was obtained by intercritical annealing at 650

for 1 h after

hot-rolling and cold-rolling of a medium-Mn TRIP steel, i.e. 0.12C-10.16Mn-1.87Al steel. The micromechanical behaviors of the used steel were investigated during tensile tests at 100, 25 and -50

, respectively, through in situ synchrotron-based 17

ACCEPTED MANUSCRIPT HE-XRD experiments. The results indicated that load partitioning took place among the constitute phases immediately after yielding. And then, the temperature-dependent stability of metastable austenite plays an important effect in the micromechanical , due to high stability, the

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behaviors at the above-mentioned temperatures. At 100

limited amount of austenite transformed to martensite, resulting in the lack of work-hardening capability. The micromechanical behavior of the used steel was just

, with the decreasing stability of retained austenite, the deformation-induced

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-50

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like to that of ultrafine-grained steels, without remarkable plastic strain. At 25 or

martensitic transformation could be triggered, leading to the improvement in the work-hardening capability. Thus, the investigated steel demonstrated pronounced plastic deformation at 25 and -50

. The austenite to martensite transformation took

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place firstly once Lüders band propagating, and lower deformation temperature promotes more volume fraction of transformed martensite while Lüders band propagation. The stability of metastable austenite in the sample deformed at 25

–was

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higher than that at -50 ℃, resulting in the stepwise transformation kinetics in the

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former condition rather than the gradual transformation at lower temperature during the stage of uniform plastic deformation. Such phenomenon resulted in the relatively lower tensile strength and much larger total elongation, as well as the best product of tensile strength and total elongation of 51.6 GPa% for the sample deformed at 25 ℃. Moreover, the stepwise phase transformations characterized by the abrupt change in volume fraction correlate with stepwise peak broadening of austenite, which are caused by PLC band propagation. While because of more intense TRIP effect at -50 ℃, 18

ACCEPTED MANUSCRIPT PLC band propagation was restrained, leading to the gradual deformation-induced transformation from austenite to martensite. Acknowledgements

(Grant No.s 51471032

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This work was supported by National Natural Science Foundation of China (NSFC) 51231002

51527801), the Fundamental Research Funds for

the Central Universities (Grant No.s 06111020 and 06111040) and the State Key

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Laboratory for Advanced Metals and Materials (Grant No.s 2016Z-01

2016Z-12

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2016Z-19). The use of the Advanced Photon Source was supported by the US Department of Energy, Office of Science, Office of Basic Energy Sciences, under Contract No. DE-AC02-06CH11357. References

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[1] O. Bouaziz, H. Zurob, M. Huang, Driving force and logic of development of advanced high strength steels for automotive applications, Steel Res. Int. 84 (2013) 937-947.

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[2] S.H. Kim, H. Kim, N.J. Kim, Brittle intermetallic compound makes ultrastrong

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ACCEPTED MANUSCRIPT Table 1 Chemical composition of the experimental Fe-0.1C-10Mn-2Al steel (mass%).

Table 2 The mechanical properties of the Fe-0.1C-10Mn-2Al steel at different

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temperatures.

25

ACCEPTED MANUSCRIPT Figure Captions: Fig.1 Schematic of the experimental set-up and notations used for in situ HE-XRD

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experiments under tensile load.

Fig.2 SEM micrograph of sample intercritically annealed at 650

strain curves of medium-Mn TRIP steels during in situ

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Fig.3 Engineering stress

for 1h.

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loading for HE-XRD experiments at temperatures of 100, 25 and -50 ℃, respectively.

Fig. 4 Selected 2D X-ray diffraction images along the full azimuthal angle η (0°-360°) for the medium-MIP steels: (a) as-received, (b) 100 ℃ at applied strain of 0.212, (c)

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25 ℃ at applied strain of 0.454 and (d) -50 ℃ at applied strain of 0.346, respectively. Noted that η =0° and 180° correspond to the transverse direction (TD), while η =90°

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and 270° are approximately along the loading direction (LD).

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Fig. 5 Evolution of austenite volume fraction as a function of applied strain for the medium-Mn TRIP steels.

Fig. 6 The α-211 and γ-311 diffraction patterns as a function of applied load at (a) 100℃, (b) 25℃ and (c) -50℃ for medium-Mn TRIP steels.

Fig.7 (a) The {211} diffraction peaks of ferrite and martensite deformed at applied 26

ACCEPTED MANUSCRIPT stress of 950 MPa; (b) the {211} diffraction peaks of ferrite and martensite deformed at applied stress of 1130 MPa; (c) d-spacing of α'-211 variation from the loading

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direction (LD) to transverse direction (TD).

Fig. 8 Lattice strains along the loading direction as a function of applied stress (a, c, e) and strain (b, d, f) for reflections from ferrite, austenite and martensite of the (c, d), and -50

(e, f).

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(a, b), 25

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medium-Mn TRIP steels at 100

Fig.9 Peak broadening along the loading direction from reflections of ferrite, austenite and martensite phases as a function of applied strain during tensile loading for (a) and (c) -50 , respectively.

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100 , (b) 25

Fig.10 Propagation of Lüders band revealed by X-ray gauge-section-scanning at different applied strain levels at 25

: (a) revealed by the change of austenite volume

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fraction; (b) revealed by the change of peak broadening of γ-311; (c) revealed by the

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development of lattice strains. The arrows point to the direction of Lüders band propagation.

Fig. 11 The spatially-resolved distribution of austenite volume fraction and peak broadening for the γ-311 reflection along the sample gauge length at two different applied plastic strains (ε=0.21 and 0.29).

27

ACCEPTED MANUSCRIPT Table Caption

Table 1 Chemical composition of the experimental Fe-0.1C-10Mn-2Al steel (mass%). Mn

Al

Si

Fe

0.12

10.16

1.87

0.05

Bal

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Table 2 The mechanical properties of the Fe-0.1C-10Mn-2Al steel at different temperatures. UTS (MPa)

TE (%)

100

940

940

21.2

25

960

1137

-50

976

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45.4

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UTS×TE (GPa%)

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YS (MPa)

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Temperature

34.6

20.0

51.6 42.1

ACCEPTED MANUSCRIPT Figure Captions

Fig.1 Schematic of the experimental set-up and notations used for in situ HE-XRD

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experiments under tensile load.

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Fig.2 SEM micrograph of sample intercritically annealed at 650

for 1h.

ACCEPTED MANUSCRIPT Fig.3 Engineering stress–strain curves of medium-Mn TRIP steels during in situ

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loading for HE-XRD experiments at temperatures of 100, 25 and -50 ℃, respectively.

Fig.4 Selected 2D X-ray diffraction images along the full azimuthal angle η (0°-360°)

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for the medium-MIP steels: (a) as-received, (b) 100 ℃ at applied strain of 0.212, (c) 25 ℃ at applied strain of 0.454 and (d) -50 ℃ at applied strain of 0.346, respectively. Noted that η =0° and 180° correspond to the transverse direction (TD), while η =90°

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and 270° are approximately along the loading direction (LD).

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ACCEPTED MANUSCRIPT Fig.5 Evolution of austenite volume fraction as a function of applied strain for the

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medium-Mn TRIP steels.

Fig.6 The α-211 and γ-311diffraction patterns as a function of applied load at (a) and (c) -50

for medium-Mn TRIP steels.

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100 , (b) 25

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ACCEPTED MANUSCRIPT Fig.7 (a) The {211} diffraction peaks of ferrite and martensite deformed at applied stress of 950 MPa; (b) the {211} diffraction peaks of ferrite and martensite deformed at applied stress of 1130 MPa; (c) d-spacing of α'-211 variation from the loading

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direction (LD) to transverse direction (TD).

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ACCEPTED MANUSCRIPT Fig. 8 Lattice strains along the loading direction as a function of applied stress (a, c, e) and strain (b, d, f) for reflections from ferrite, austenite and martensite of the (a, b), 25

(c, d), and -50

(e, f).

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medium-Mn TRIP steels at 100

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ACCEPTED MANUSCRIPT Fig.9 Peak broadening along the loading direction from reflections of ferrite, austenite and martensite phases as a function of applied strain during tensile loading for (a) and (c) -50 , respectively.

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100 , (b) 25

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Fig.10 Propagation of Lüders band revealed by X-ray gauge-section-scanning at different applied strains at 25

: (a) revealed by the change of austenite volume

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fraction; (b) revealed by the change of peak broadening of γ-311; (c) revealed by the development of lattice strains. The arrows point to the direction of Lüders band

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propagation.

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Fig. 11 The spatially-resolved distribution of austenite volume fraction and peak broadening for the γ-311 reflection along the sample gauge length at two different applied plastic strains (ε=0.21 and 0.29).

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