Suppressed deformation instability in the twinning-induced plasticity steel-cored three-layer steel sheet

Acta Materialia 147 (2018) 304e312

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Suppressed deformation instability in the twinning-induced plasticity steel-cored three-layer steel sheet Jung Gi Kim a, Seung Mi Baek a, Hak Hyeon Lee a, Kwang-Geun Chin b, Sunghak Lee a, Hyoung Seop Kim a, c, * a b c

Department of Materials Science and Engineering, Pohang University of Science and Technology (POSTECH), Pohang, 37673, Republic of Korea Graduate Institute of Ferrous Technology, Pohang University of Science and Technology (POSTECH), Pohang, 37673, Republic of Korea Center for High Entropy Alloys, Pohang University of Science and Technology (POSTECH), Pohang, 37673, Republic of Korea

a r t i c l e i n f o

a b s t r a c t

Article history: Received 1 October 2017 Received in revised form 11 December 2017 Accepted 17 January 2018 Available online 3 February 2018

Deformation instabilities (i.e., yield point phenomenon (YPP) and serrated flow) were investigated for twinning-induced plasticity (TWIP) steel-core with low-carbon (LC) steel outer layer sheets. During tensile tests, both the YPP and serrated flow were suppressed as the volume fraction of the TWIP steelcore decreased. Electron backscattering diffraction analysis indicated that a strain gradient was induced, and that geometrically necessary dislocations (GNDs) accumulated at the TWIP-LC steel interface. With increase in the dislocation density by plastic deformation, the deformation instabilities in the TWIPcored layer steel sheets were suppressed due to the reduced probability of interaction of dislocation with carbon-solutes. This result indicates that the strain gradient at the interface of the layered sheets not only enhances the mechanical property, but also suppresses deformation instability. © 2018 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

Keywords: TWIP steel Architectured materials Serration Yield point phenomena Interface

1. Introduction Most monolithic metallic materials present a trade-off between strength and elongation [1]. This trade-off relationship indicates that monolithic materials have a limitation in achieving high strength with great elongation. In steel products, the strengthelongation trade-off is depicted as a ‘banana curve’ (see Fig. 1) [2]. Many researchers have tried to fill the empty area of the strength-elongation window in the Ashby diagram by developing new alloy systems [3]. Recently, architectured materials have received much attention in attempts to fill the traditional empty zone of the strength-elongation window with new materials. The concept of architectured material proposed by Ashby is the combination of two or more materials, which results in expanding the multi-functional availability of materials by designing the components, volume fraction, configuration, and connection [3,4]. By using various design architectures, the architectured materials can represent specific mechanical properties not only by following the

* Corresponding author. Department of Materials Science and Engineering, Pohang University of Science and Technology (POSTECH), Pohang, 37673, Republic of Korea. E-mail address: [email protected] (H.S. Kim). https://doi.org/10.1016/j.actamat.2018.01.042 1359-6454/© 2018 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

rule-of-mixtures [5], but also by exhibiting ‘synergetic strengthening’ [6]. Zhu et al. reported that gradient-structured material (that with a strain gradient) originates from stress incompatibility at the interface between fine and coarse grains [7]. This induces extra strain near the interface region. To accommodate the strain incompatibility at the grain interface, geometrically necessary dislocations (GNDs) are generated in the interface between fine and coarse grains [8]. In general, the dislocation density of a material has a square-root relationship with the strength (i.e., strength f √(dislocation density)), and the increased GNDs in the gradient material enhance the strength as well. By this is meant that the interface in the gradient material induces extra strengthening in addition to the conventional deformation strengthening. This extra strengthening occurs not only in the gradient materials, but also in the lamellar or laminate structure [9,10]. Recently, it was reported for a copper and brass multilayer laminated structure processed by accumulative roll bonding and annealing, that the amount of GND accumulation during tensile testing depended on the thickness of each layer [11]. The GND density increased as the layer thickness decreased, and the tensile strength of the laminated material was enhanced by the GNDs [11]. Thus, the copper/brass laminated material exhibited extra

J.G. Kim et al. / Acta Materialia 147 (2018) 304e312

Ultra High Strength Steels

60

High Strength Steels

50

MILD

Generation AHSS

TWCLS

IS

I IF-HS

20

30

IF

40

2

10

Elongation, %

Low Strength Steels

MART

0

300

600

900

1200

1500

Tensile Strength, MPa Fig. 1. Strength-elongation trade-off relationship in steel product which called ‘banana curve’. The 2nd generation advanced high strength steel (AHSS) is the representative product to fill up the vacant area (Gray-marked area) of the banana curve [2]. The combination of TWIP steel e mild steel covers the vacant area of the banana curve.

strengthening due to the strain gradient at the interface. Such an advantage in laminated structures is enough to gain the attention of material engineers and scientists, and this strategy can be applied to steel products. In fact, an approach similar to gradient architectures has already been created for steel products because the lifetime of machining tools is such an important issue. Surface treatment processes (i.e., accumulative roll bonding [12], surface mechanical attrition treatment [13], and carburization [14]) induce not only grain refinement in the surface, but also gradient microstructures. Layered architectures have been applied primarily to sheets of high strength steels (HSSs). Layered sheets of HSS combined with another steel have greater strength. For example, Bouaziz et al. reported the enhanced mechanical property (900 MPa tensile strength, with 35% elongation) of twinning-induced plasticity (TWIP)-martensitic bi-layer steel sheet [15]. The strength followed the simple rule-of-mixture relation of the constituent materials. Koseki et al. studied multilayered steel constructed with martensitic and stainless steels [16e18], and achieved more than 1200 MPa tensile strength and 20e25% tensile elongation. Moreover, the multi-layered steel sheet had a moderate forming limit similar to those of commercial dual-phase steels and transformation-induced plasticity steels [19]. In addition, the strain partitioning and interfacial hardening in the HSS-based layered materials increased due to the great difference in strength between the two parent materials. While strain partitioning occurs in layered materials, the hard phase is less deformed than the soft phase, and extra-strengthening occurs from the soft phase. These micro-mechanical phenomena can be linked to the strengthening of multiphase or lamellar structures, which give extra-strength to the material [20]. Therefore, understanding multilayered steel is important not only for designing the properties of steels, but also for unveiling the micromechanics of the interface. From review of the recent status in the field of architectured materials, particularly in the authors' previous research, the development of TWIP-cored three-layer steel (TWCLS) sheet is represented in Fig. 2 [21]. It should be noted that both TWIP and low carbon (LC) steel monolithic layers have deformation instabilities such as yield point phenomena (YPP) and dynamic strain aging (DSA) during plastic deformation [22e25]. These deformation instabilities induce strain localization and have a negative effect on the steel product. It should be noted that the YPP results from the interaction between dislocations and interstitial solutes [26,27]. When the carbon solute atoms are assembled with

305

dislocations, the dislocations cannot escape from this atmosphere (Cottrell atmosphere) with a small force. To break the Cottrell atmosphere, an additional force larger than the energy barrier for dislocation gliding should be imposed. After the dislocation gliding occurs, the energy barrier is reduced, and the surplus force flows rapidly after yielding [27]. To suppress the YPP in steel products, the interaction between the interstitial solute and dislocations should be reduced, and some approaches (i.e., increasing the dislocation density or reducing the interstitial solutes) are well established. Keh et al. reported that the YPP of quenched LC steel was suppressed. In the quenched LC steel, high dislocation density existed, and the probability of interaction between the carbon-solute and dislocations was reduced [28]. Similarly, DSA arises by interaction between dislocations and solute clusters. To explain this phenomenon, Mulford et al. investigated solute atom diffusion along the core of mobile dislocations [29]. Their model is generally accepted because it has been verified with several experimental results [30,31]. Thus, it is clear that the interaction between solute atoms and dislocations should be reduced to inhibit DSA during plastic deformation. One efficient way is adding aluminum to the TWIP steel, which results in suppressed carbon-manganese clustering in TWIP steels [32]. The interaction between solute clusters and mobile dislocations is reduced and the DSA is decreased during plastic deformation, as the amount of aluminum increases. Another way is increasing the dislocation density, which reduces the DSA stress contribution in the TWIP steel. Because the origin of DSA is related to the YPP, the DSA can be treated as the YPP at the location of high strain. Kim et al. quantified the effect of DSA on strain hardening by measuring the yield stress increment during the loading-unloading-agingreloading procedure [33]. This result implies that if the YPP at the high strain is suppressed, the DSA will be absent. Thus, similarly to YPP reduction, dislocation density enhancement could be an appropriate method to inhibit the occurrence of DSA. Therefore, dislocation density enhancement is one efficient way to reduce deformation instabilities, and we believed that the accumulation of GND in the architectured material would have a similar effect on TWCLS sheets. In this study, the deformation instability of the TWCLS sheets was investigated. The strain and strain-rate distributions of TWCLS sheets were measured using the digital image correlation (DIC) method. The YPP of TWCLS sheets was examined in terms of the volume fraction of the parent materials to check the reduction of YPP as the volume fractions of TWIP and LC steels varied. The serrated flow of each specimen was observed in the late strain region of the stress-strain curves. To clarify the strain gradient in a three-layer steel sheet, finite element method (FEM) analysis was conducted to show the strain distribution in the near-interface region. The GND distribution was quantified using electron backscattering diffraction (EBSD) analysis. To check the relationship between YPP and dislocation density, the overall dislocation density was calculated by using the combination of X-ray diffraction (XRD) and peak profile analyses. 2. Experimental procedure The chemical compositions of the TWIP and LC steels are Fe15Mn-1.2Al-0.6C and Fe-0.2Mn-0.04Al-0.03C, respectively. TWCLS sheets were manufactured at POSCO (South Korea), and Fig. 2 shows the sandwich structure of the layered sheets. The volume fractions of the three-layer steel sheets were controlled by changing the stacking ratio of the parent materials. Table 1 represents the TWIP and LC steel volume fractions. The name of the TWCLS sheets was designated according to the stacking ratio (i.e., CLAD 1/2/1 for a 25% LC steel-sheath, 50% TWIP steel-core, and 25% LC steel-sheath

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Fig. 2. Optical microscope and EBSD results of the TWCLS sheet. (a) Optical microscopy of the CLAD 1/1/1 TWCLS sheet; and (b) EBSD image measured at the TWIP-LC interface.

Table 1 The volume fractions (%) of the TWCLS sheets. Specimen

LC

TWIP

LC

CLAD 1/6/1 CLAD 1/2/1 CLAD 1/1/1

12.5 25 33

75 50 33

12.5 25 33

TWCLS sheet). The stacked steel sheets were homogenized at 1200  C for one hour and hot rolled from 40 to 2.5 mm thickness within a 900e1100  C temperature regime. After hot rolling, the sheets were cold-rolled from 2.5 to 1 mm thickness and annealed at 820  C for 30 s. The tensile specimens were prepared with a 5 mm gage length plate-type sub-sized specimen, and the tests were done using a universal testing machine (Instron 1361, Instron Corp., Canton, MA, USA) at room temperature with 1  10 3 s 1 quasi-static strain rate. The strain and strain rate distributions of the tensile specimens were measured using the DIC (ARAMIS v 6.1, GOM Optical Measuring Techniques, Germany) method with white and black speckles on the surface of tensile specimens [34]. The changes in the strain gradient at the TWIP-LC steel interface were estimated using FEM (ABAQUS 6.9/EF-2, Dassault Systems, France). To ensure the simulation result, each layer has at least 5 meshes along the thickness direction and the fully integrated hexahedral elements (C3D8) were used. To correlate the FEM results with the experimental results, the tensile simulation conditions were made the same as those for the tensile tests, in which the tensile specimen (5 mm gage-length plate-type sub-sized) was deformed at room temperature, with a quasi-static strain rate of 1  10 3 s 1. The material property of the TWIP steel-core and LC steel-sheath were incorporated into the tensile specimens. To consider the anisotropic properties of the TWIP steel-core and LC steel-sheath layers, R-values at 0 (R0), 45 (R45), and 90 (R90) degrees to the rolling direction were included, as defined in Table 2 [35]. To implement anisotropic properties into the ABAQUS software, the deformation state of the tensile specimens was treated as a planar anisotropy condition. Each anisotropic potential value was calculated based on the Hill's 1948 anisotropy criterion [36]. To clarify the GND accumulation in the TWIP-LC interface

Table 2 The R0, R45, and R90 of the TWIP and LC steels. Specimen

R0

R45

R90

TWIP steel LC steel

0.68 0.96

0.82 0.85

1.03 1.08

region, GND-distributions in the TWIP steel-core and LC steelsheath were quantified using EBSD analysis (XL-30S FEM, Philips Electron Optics, The Netherlands). The specimens were polished using mechanical and colloidal processes. The step size was 65 nm for each specimen with a 0.09 minimum confidence index (CI) value to obtain high-quality results. The XRD (D/MAX-2500, RIGAKU, Japan) patterns of the TWCLS sheets were measured from 30 to 110 . To observe the TWIP steel-core, the XRD was performed at the side surface of the 5% deformed tensile specimen. The dislocation density of the TWIP steel-core was calculated using the convolutional multiple whole profile (CMWP) method [37]. 3. Results 3.1. Tensile test Fig. 3(a) shows the engineering stress-strain curves of the TWCLS sheets with engineering strain e < 0.1. The yield point elongation (YPE%) of the TWCLS sheets decreased as the volume fraction of TWIP steel-core decreased, while both the monolithic TWIP steel and LC steel have ~1 and ~5% YPE%, respectively. Fig. 3(b) shows the true stress-strain curves of the TWCLS sheets in 0.2 < true strain ε < 0.5 when observing serrations. Serrated flow was observed in the monolithic TWIP steel while the TWCLS sheets had smooth flow without serration. The previous DIC result shows that the Portevin-Le Chatelier (PLC) band propagation is suppressed in the TWCLS sheet while the PLC band propagation is clearly observed in the monolithic TWIP steel [35]. For detailed analysis of the YPE% variations in the TWCLS sheets, the YPE% changes as a function of the volume fraction of the TWIP steel-core are shown in Fig. 4. YPE% decreased from 1 to 0% as the volume fraction of the TWIP steel-core decreased, while CLAD 1/1/1 showed 0 YPE%. To investigate the Lüders band propagation in the TWCLS sheets, strain and strain-rate distributions were measured using the DIC method [37]. Fig. 5 represents the strain and strainrate distributions of the TWCLS sheets during the tensile test, in which the DIC was measured on the side surfaces. The strain and strain-rate distributions of CLAD 1/6/1 (Fig. 5(a)), which had the largest TWIP steel-core volume fraction, showed strain-rate localization. This strain-rate localization originated from Lüders band propagation (4e16 s), which induces strain localization during tensile deformation. The DIC measurement results for CLAD 1/2/1 (Fig. 5(b)) shows the Lüders band propagation from 4 to 12 s. Its intensity and holding time are lower than those in CLAD 1/6/1, and strain localization is reduced due to the weakened Lüders band. The results for the CLAD 1/1/1 sheets in Fig. 5(c) indicate that the Lüders band propagation was suppressed and that the tensile specimen deformed uniformly. The strain and strain rate distributions of the

J.G. Kim et al. / Acta Materialia 147 (2018) 304e312

(a)

1400

700 600 500 400 300 TWIP CLAD1/6/1 CLAD1/2/1 CLAD1/1/1 Low Carbon

200 100 0 0.00

0.02

0.04

0.06

0.08

0.10

Engineering Strain

Engineering Stress, MPa

Engineering Stress, MPa

800

307

(b)

1200 1000

TWIP CLAD1/6/1 CLAD1/2/1 CLAD1/1/1 Low Carbon

800 600 400 200 0.20

0.25

0.30

0.35

0.40

0.45

0.50

Engineering Strain

Fig. 3. Stress-strain curves of the TWIP, LC, and TWCLS sheets after tensile tests. (a) Initial strain region; and (b) Lateral strain region.

3.2. Finite element analysis

Yield Point Elongation, %

8 7 Single Layer TWCLS Sheet

6 5 4 3 2 1 0

0

20

40

60

80

100

TWIP Steel Volume Fraction, % Fig. 4. Plots of the YPE% of TWCLS sheets as the volume fraction of TWIP steel-core changes.

TWCLS sheets show that the propagation time of the Lüders bands decreased as the volume fraction of the TWIP steel-core decreased. The origin of the reduced Lüders band propagation time will be discussed later. These results imply that strain localization was suppressed in the TWCLS sheet and that homogeneous deformation occurred during plastic deformation.

Fig. 6 represents the comparison results of the equivalent plastic strain distributions of the 5% deformed CLAD 1/2/1 sheet from the FEM simulation and the DIC result. Fig. 6(a) and (b) show the equivalent plastic strain distribution map from the FEM and DIC method, respectively. In the FEM simulation result, strain partitioning between TWIP steel-core and LC steel-sheath was observed. This strain partitioning originated from the strength difference between the hard TWIP steel-core and soft LC steel-sheath. However, with the DIC method, strain partitioning was not clearly observed due to the gradient contour. For detailed analysis, the equivalent plastic strain profiles were extracted from the marked paths and plotted. From Fig. 6(c), we found that the clear strain partitioning in the DIC method occurs similarly to that in the FEM simulation result. This result shows that the reliable strain distributions of TWCLS sheets from the FEM simulation can be correlated with the experimental result. Fig. 7(a) shows the thickness strain distribution of the deformed CLAD 1/2/1 sheet. Because of the high R-value of LC steel-sheath, the LC steel sheath has a high thickness reduction resistance, and lower thickness strain is observed at the edge (side) of the tensile sample than that in the TWIP steel-core. In the central region, however, the thickness strain distribution is inverted due to the large strength difference between the TWIP steel-core and LC steel-

Fig. 5. Strain and strain rate distributions of TWCLS sheets during tensile tests. (a) CLAD 1/6/1; (b) CLAD 1/2/1; and (c) CLAD 1/1/1. The strain and strain rate distribution results were extracted from initiation to annihilation of Lüders bands with a 4 s time interval.

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Fig. 6. Equivalent plastic strain distributions of 5% tensile elongated CLAD 1/2/1 TWCLS sheet from (a) FEM simulation and (b) DIC method. (c) The equivalent plastic strain profiles along the ‘PATH’ in (a) and (b).

Fig. 7. FEM simulation results of the (a) thickness and (b) width strain distributions of 5% tensile elongation CLAD 1/2/1 TWCLS sheet.

sheath. In contrast, as represented in Fig. 7(b), the width strain distribution of the CLAD 1/2/1 is opposite to the thickness strain distribution. The high R-value of the LC steel-sheath induces larger width strain than the TWIP steel-core, and results in substantial

shrinkage in the width direction. As a result, the stress state of the tensile deformed TWCLS sheet not only consisted of a uniaxial tensile deformation component, but also included triaxial stress components. This means that both the mechanical and the anisotropic property differences between TWIP steel-core and LC steelsheath should be considered when analyzing evolution of plastic strain in TWCLS sheets. These differences in mechanical and anisotropic properties in the TWIP steel-core and LC steel-sheath, induce a strain gradient at the TWIP-LC interface though a constitutive layer interaction at the TWIP-LC interface, which was not considered in this simulation [7]. Fig. 8 shows the strain and strain gradient profiles of the 5% deformed TWCLS sheets. Each profile was extracted from the side edge of the specimens, and the results show that a strong strain gradient was induced at the TWIP-LC interface. The strain gradient of the TWCLS sheet was not only observed in the LC steel-sheath but was also effective in the TWIP steel-core. The strain gradient was most extreme at the TWIP-LC interface of CLAD 1/6/1. This additional strain gradient originated from the heterogeneous strain profile along the thickness direction of the monolithic materials during tensile tests. Fig. 9 shows the equivalent plastic strain profiles along the thickness direction of the 5% tensile deformed monolithic TWIP steel, and the profile revealed that heterogeneous strain distribution occurred along the thickness direction during

J.G. Kim et al. / Acta Materialia 147 (2018) 304e312

309

0.044

CLAD 1/6/1 CLAD 1/2/1 CLAD 1/1/1

0.043

Loading

Equivalent Plastic Strain

(a)

0.045

0.042

0.041

-1

CLAD 1/6/1 CLAD 1/2/1 CLAD 1/1/1

0.007 0.006 0.005

Loading

Strain Gradient, mm

PATH

(b)

0.008

0.004 0.003 0.002 0.001 0.000 0.0

0.2

0.4

0.6

0.8

1.0

Path Distance, mm Fig. 8. The equivalent plastic strain and strain gradient profiles of the 5% tensile elongated TWCLS sheets. (c) The strain and strain gradient are extracted at the edge side, and along the thickness direction of TWCLS sheet.

the tensile test. This strain inhomogeneity enhances the strain gradient at the TWIP-LC interface, and this effect is stronger at the outside of the tensile sample than in the central region. Although the strain gradient was larger for CLAD 1/6/1 than for CLAD 1/1/1, the central region of the TWIP steel-core in CLAD 1/6/1 was not

Equivalent Plastic Strain

0.04385

Surface

Edge

0.04384 0.04383 0.04382 0.04381 0.04380 0.0

0.1

0.2

0.3

0.4

0.5

Distance from Surface, mm Fig. 9. Equivalent plastic strain profile along the thickness direction of the 5% tensile elongated monolithic TWIP steel.

affected by the strain gradient due to the great thickness of the TWIP steel-core.

3.3. EBSD analysis The plastic strain gradient was linked to the stored excess of GNDs using Nye's dislocation tensor [38,39]. The EBSD analysis is one efficient method for quantifying the local rotation gradients, and these local gradients can be transformed to the GND density [40]. In this study, the GND distributions of the TWCLS sheets were measured using EBSD analysis along the thickness direction [11]. Fig. 10(a) shows the GND distribution map of the CLAD 1/2/1 sheet and large GND accumulation is observed at the interface (marked with arrows). Fig. 10(b) and (c) show the GND density distributions at the TWIP steel-interface and center, respectively. These diagrams show that the average GND density of TWIP steel-interface (2.63  1014 m 2) is higher than that in the center region (2.58  1014 m 2). Fig. 10(d) and (e) represent the GND density distributions at the LC steel-interface and edge, respectively. As with the result of the TWIP steel-core, the average GND density of the LC steel-interface (1.78  1014 m 2) is larger than that at the edge region (1.70  1014 m 2). These results coincided with the strain gradient plots (as in Fig. 8(b)), indicating that the high strain gradient at the interface induces GND accumulation in the TWCLS sheets.

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Fig. 10. (a) GNDs distribution maps of the 5% tensile elongated CLAD 1/2/1 TWCLS sheet and the GND density histograms from the (b) TWIP steel-interface, (c) TWIP steel-center, (d) LC steel-interface and (e) LC steel-edge.

4. Discussion 4.1. GND accumulation in the interface of TWCLS sheets The EBSD analysis and FEM simulation results reveal that the TWIP-LC interface induces GND accumulation in the TWCLS sheet and strain partitioning occurs due to the strength difference between TWIP steel-core and LC steel-sheath. Fig. 11 is a schematic drawing of the deformation mechanism at the TWIP-LC interface: (i) In Stage A, both TWIP steel-core and LC steel-sheath are in elastic deformation. (ii) In Stage B, the LC steel-sheath starts plastic deformation while the TWIP steel-core is under elastic deformation. Because of the constraint region near the TWIP-LC interface, plastic strain gradient occurs at the LC steel-sheath. (iii) In Stage C, both TWIP steel-core and LC steel-sheath are in plastic

deformation. As with Stage B, different properties of the parent materials induced a strain gradient in both the TWIP steel-core and the LC steel-sheath. Based on the TWIP-LC interfacial deformation, additional GNDs evolved at the TWIP-LC interface as represented in Fig. 10. Based on this result, we can estimate that the overall dislocation density of the TWCLS sheet can be increased by additional GNDs. However, the current EBSD analysis result has a limitation to get a reliable dislocation density due to the relatively large step size. According to the previous research, the GND density from EBSD analysis has an inversely proportional relation with the step size [41]. Therefore, another dislocation density quantification method is required to prove the increased overall dislocation density in the TWCLS sheet and the X-ray peak profile analysis is an efficient method to quantify overall dislocation density.

Fig. 11. Schematic drawing on the elasto-plastic deformation mechanism at the TWIP-LC interface. (a) STAGE A; (b) STAGE B; (c) STAGE C.

J.G. Kim et al. / Acta Materialia 147 (2018) 304e312

Fig. 12 plots the relationship between YPE% and dislocation density of the TWIP steel-core (rTWIP-core) in the 5% tensile deformed TWCLS sheet. As a result, the YPE% is decreased as the dislocation density increases and the dislocation density of the TWCLS sheet is larger than that in the monolithic TWIP steel (rMonolithic-TWIP). The dislocation density of the TWCLS sheet is the largest at CLAD 1/1/1 sample while the CLAD 1/6/1 and 1/2/1 samples show relatively small dislocation densities. Such a different dislocation density evolution in each TWCLS sheet can be explained by the strain gradient distributions in Fig. 8 and the different layer configuration. Because of the large layer thickness of the CLAD 1/6/1 and 1/2/1 samples, the strain gradient cannot affect the center region and this limitation results in the less dislocation density enhancement.

4.2. Deformation instability suppression in TWCLS sheets Moon et al. reported that the band propagation velocity (VB) is proportional to the dislocation density [42]. If the GNDs are accumulated at the TWIP-LC interface, the overall dislocation density of the TWCLS will increase. The elevated dislocation density increases VB in TWCLS sheet and decreases the Lüders band propagation time. However, the strain gradient of TWCLS sheet becomes ineffective when the layer thickness is too great, as shown in Fig. 8(b). This means that the GND accumulation in the central region of the TWIP steel-core will decrease as the thickness of the TWIP steelcore increases. Because of less GND accumulation in the central region of CLAD 1/6/1, the Lüders band propagation time also increased with layer thickness and Lüders band propagation was observed. The stress-strain curve in Fig. 3(b) revealed that the TWCLS sheets suppress the serrated flow. This suppression can be explained by the relationship between YPP and the triaxial deformation mode during tensile tests. First, as mentioned in the introduction, the DSA of TWIP steel is theoretically similar to the YPP [33]. This means if the interaction between dislocation and solute atoms is reduced, the DSA of the TWCLS will be suppressed. Similar to quenching in LC steel [28], the accumulated GNDs at the TWIP-LC interface results in high dislocation density in TWCLS sheets and this high dislocation density reduces the probability of interaction with the solute atoms. Second, several previous reports show that biaxial or triaxial deformation reduces serration at the large strain [43,44]. Because of the different mechanical and anisotropic properties between the TWIP steel-core and the LC steel-sheath, the strain distribution in tensile deformed TWCLS

Fig. 12. The relationship between TWIP steel-core (rTWIP-core)/monolithic TWIP steel (rMonolithic-TWIP) and YPE%.

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sheets is not the same as for a uniaxial tensile state. Both the FEM simulation and DIC results reveal that the strain partitioning and triaxial stress state are induced at the TWIP-LC interface during tensile tests. These triaxial stresses of the TWCLS sheets can be linked with previous results that triaxial deformation reduces serrated flow [44,45]. 5. Conclusions In summary, the deformation instability suppression in TWCLS sheet is governed from the TWIP-LC interface. During tensile tests, plastic deformation started from the soft LC steel-sheath, and the EBSD analysis result revealed that GND accumulation occurs to cover the strain gradient at the TWIP-LC interface. The accumulated GNDs increase the overall dislocation density in TWCLS sheet and reduce the interaction between dislocation and solute atoms. Moreover, the FEM and DIC results show that the strain partitioning and triaxial stress state occur at the TWIP-LC interface in the tensile deformed TWCLS sheet due to differences in the mechanical and anisotropic properties of the parent materials. The reduction of interaction and triaxial stress state in the TWCLS sheets result in YPP and DSA suppression. These results show that recent properties of these architectured materials not only follow the theoretical rule-of-mixtures, but also have potential to overcome the limitations of previous materials. Acknowledgements This study was supported by POSCO (20148053). This work was supported by the National Research Foundation of Korea (NRF) grant funded by the Ministry of Science and ICT (MSIP) (No. 2014R1A2A1A10051322). Also, this study was supported by Brain Korea 21 PLUS project for Center for Creative Industrial Materials (F16SN25D1706). One of the author (JGK) acknowledges the support from the Global Ph.D. Fellowship funded by the National Research Foundation of Korea (No. 2013H1A1032565). References [1] Y. Wei, Y. Li, L. Zhu, Y. Liu, X. Lei, G. Wang, Y. Wu, Z. Mi, J. Liu, H. Wang, H. Gao, Evading the strength-ductility trade-off dilemma in steel through gradient hierarchical nanotwins, Nat. Commun. 5 (2014) 3580. [2] O. Bouaziz, S. Allain, C.P. Scott, P. Cugy, D. Barbier, High manganese austenitic twinning induced plasticity steels: a review of the microstructure properties relationships, Curr. Opin. Solid State Mater. Sci. 15 (2011) 141e168. [3] M. Ashby, Designing architectured materials, Scr. Mater. 68 (2013) 4e7.
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