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Prediction of tensile properties of intercritically annealed Al-containing 0.19C–4.5Mn (wt%) TRIP steels
Materials and Design 97 (2016) 138–146
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Prediction of tensile properties of intercritically annealed Al-containing 0.19C–4.5Mn (wt%) TRIP steels Singon Kang a, John G. Speer a, Daniel Krizan b, David K. Matlock a, Emmanuel De Moor a,⁎ a b
Advanced Steel Processing and Products Research Center, Colorado School of Mines, Golden, CO 80401, USA Research and Development Department Business Unit Coil, Voestalpine Steel Division GmbH, Voestalpine-Str. 3, Linz A-4031, Austria
a r t i c l e
i n f o
Article history: Received 3 November 2015 Received in revised form 8 February 2016 Accepted 15 February 2016 Available online 17 February 2016 Keywords: Transformation-induced plasticity Retained austenite Composite model Intercritical annealing Alloy partitioning
a b s t r a c t Tensile properties were predicted for intercritically annealed 4.5 wt% Mn steels with varying Al additions up to 1.3 wt%. Intercritically annealed phase fractions and retained austenite chemical compositions were obtained using a retained austenite prediction model assuming ortho-equilibrium alloy partitioning. The martensite transformation kinetics during tensile deformation were also estimated from the predicted retained austenite chemical compositions. Predicted phase fractions and martensite transformation kinetics were incorporated in a composite model for multi-component systems. Predicted tensile strength and uniform elongation values were compared to the experimental data obtained from the steels annealed at temperatures ranging from 550 to 725 °C for 1 h followed by water quenching. © 2016 Elsevier Ltd. All rights reserved.
1. Introduction Recent developments in “third generation” advanced high strength steels (3G-AHSS) for automotive applications aim to achieve high strength and ductility to meet the needs for fuel efficiency and passenger safety. Phase transformation from austenite to martensite during tensile deformation, i.e. transformation-induced plasticity (TRIP), is one of the prominent strategies to obtain desired tensile properties for 3G-AHSS [1]. Through the strain-induced or stress-assisted transformation of austenite to the harder martensite phase, exceptional strain hardening occurs during tensile deformation, improving formability of these high strength steels [2]. Since the work by Miller [3] on a 0.11C–5.7Mn steel, medium Mn steels have been extensively investigated [4–14]. Heat treatment of medium Mn steels in the intercritical temperature range enables partitioning of alloying elements between ferrite and austenite. Austenite enrichment by stabilizing elements such as Mn and C restricts martensite formation during quenching and increases retained austenite fractions (fγR). To obtain reasonable amounts of fγR by intercritical annealing and understand the effects of retained austenite on tensile properties, studies have been performed using medium Mn TRIP steels [7–9]. Huang et al. [8] annealed a 0.12C–5.1Mn steel using a variety of temperatures, holding times and cooling conditions. For 3 h annealing followed by a water quench, a variation in annealing temperature from 700 to ⁎ Corresponding author. E-mail address: [email protected] (E. De Moor).
http://dx.doi.org/10.1016/j.matdes.2016.02.058 0264-1275/© 2016 Elsevier Ltd. All rights reserved.
650 °C increased the fγR from about 10 to 30%, and improved the product of ultimate tensile strength (UTS) and uniform elongation (UE) values from 10,000 to 25,000 MPa·%. Kim et al. [9] reported that the UTS×UE values for 0.1 C steels containing 4, 6 or 8 wt% Mn annealed in the intercritical temperature range, increased with an increase in fγR. This confirms that tensile properties of medium Mn TRIP steels may be enhanced by increased fγR. The importance of austenite fractions has been reported by others as well, along with austenite stability [4,5,10,11]. The effect of martensite transformation behavior during tensile deformation on the tensile properties of intercritically annealed medium Mn steels has been examined e.g. [10,11]. Gibbs et al. [10] performed intercritical annealing of a 0.1C–7.1Mn–0.1Si steel in a temperature range of 575 to 675 °C for 168 h followed by water quenching. Although less austenite was stabilized by annealing at 575 °C (26.3%) compared to 650 °C (43.5%), the 650 °C sample exhibited reduced elongation values compared to the 575 °C sample. This behavior was ascribed to the rapid transformation kinetics in the 650 °C sample as more than 80% of the retained austenite transformed to martensite within a true strain range of 0.05. A similar reduction in austenite stability with increasing annealing temperature was observed by Suh et al. [11] in a 0.11C– 4.5Mn–0.5Si–2.2Al steel. In addition to the amount and stability of austenite, the fractions of other phases such as ferrite, cementite and martensite are also greatly influenced by the intercritical annealing conditions [10], and correspondingly the resulting tensile properties depend on the ferrite, cementite and martensite fractions in annealed medium Mn steels. While there are multiple approaches currently being considered to predict the properties of multi-component systems which evolve with
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imposed macroscopic strain to account for strain and stress gradients between individual constituents, a modified rule-of-mixture approach is frequently used for predictions of tensile properties and it provides an insight into design requirements for new steels with complex microstructures. Matlock and Speer [1,12] used the modified rule-of-mixture approach to propose a composite model for predictions of tensile properties incorporating martensite transformation kinetics during tensile deformation. The tensile flow curves for each phase were assumed to follow Hollomon's power law, and the strain dependent phase fractions were applied to consider the effect of martensite transformation during the tensile deformation. Predicted variations in tensile properties of multi-phase steels were provided for different combinations of phase fractions and martensite transformation kinetics. Rana et al. [13] applied the composite model to an experimental alloy having a chemical composition of 0.1C–7.1Mn–0.1Si annealed at 600, 625 and 650 °C. Their phase fractions and martensite transformation kinetics were experimentally measured [10] and model predictions were shown to directly predict observed strengths and ductilities. Lee et al. [14] predicted the tensile properties of a 0.08C–6Mn– 2.0Al–1.5Si–0.08V steel annealed at 740 °C for 3 min using a modified Estrin and Mecking constitutive model [15]. Respective constitutive models for each phase such as ferrite, austenite, and martensite were constructed using diverse experimental parameters including dislocation density, grain size, and martensite transformation kinetics. These constitutive models were combined by considering strain partitioning between the phases based on iso-work theory. These predictions successfully describe the need for significant amounts of retained austenite with controlled stability against deformation for new 3G-AHSS. However, the application of these models was limited for sample conditions with existing experimental results such as initial phase fraction and martensite transformation kinetics. In the present study, a method is proposed to predict the tensile properties of intercritically annealed multi-phase steels, which does not require pre-existing experimental data. A composite model proposed by Matlock and Speer [1,12] which incorporates predicted initial phase fraction and martensite transformation kinetics is used to obtain the tensile properties. Phase fractions and compositions for annealed samples are predicted using a model assuming ortho-equilibrium alloy partitioning during intercritical annealing [16,17]. Martensite transformation kinetics during tensile deformation is estimated based on the retained austenite chemical composition. This method is applied to three medium Mn steels containing different levels of Al with varying intercritical annealing temperature, which have different combinations of various phases and a wide range of fγR. For verification of the proposed model, predicted results are compared to experimental tensile properties obtained from equivalent alloy and heat treatment conditions. The evolution of the tensile properties with annealing temperature in both prediction and experiments are compared and discussed. 2. Experimental procedure Three 0.19C–4.5Mn (wt%) medium Mn steels containing 0, 0.5 or 1.3 wt% Al, referred to as Base, Lo-Al and Hi-Al, having the chemical compositions provided in Table 1, were designed and laboratory cast as 90 kg ingots [18]. Sample blocks having a thickness of 130 mm were reheated to 1180 °C for 2 h, hot rolled in 12 passes to a thickness of 4 mm with a finishing temperature of 900 °C, and coiled at 500 °C. These hot rolled plates had fully martensitic microstructures. In order Table 1 Chemical compositions of the Base, Lo-Al and Hi-Al alloys. wt%
C
Mn
Al
Base Lo-Al Hi-Al
0.190 0.189 0.195
4.39 4.45 4.52
0.01 0.48 1.30
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to facilitate cold rolling without crack formation, the plates were softened by a prior batch annealing at 550 °C for 16 h where the microstructure was primarily ferrite and thus alloy partitioning to austenite would be minimized. These softened plates were cold rolled in 5 passes to a final thickness of 1 mm. Tensile specimens were machined along the rolling direction with a gauge length of 50.4 mm according to the ASTM-E8 standard [19]. The specimens were isothermally heat treated at temperatures ranging from 550 to 725 °C for 1 h in a box furnace and water quenched. To avoid sample oxidation, the samples were sealed in stainless steel bags with Ti particles to getter oxygen. Tensile tests for the heat treated samples were performed at a constant strain rate of 6.25 × 10−4 s−1 at room temperature. Heat treated samples for X-ray diffraction (XRD) analysis were ground with 1200 grit SiC paper and chemically polished with a 1:50:50 mixture of 48% hydrofluoric acid, 30% hydrogen peroxide, and de-ionized water at room temperature for 180 s. A Philips X-pert diffractometer with a copper tube was operated at 45 kV and 40 mA and a 2θ scan range from 40 to 110° was employed. After stripping the Kα2diffracted peaks through X-Pert software, the intensities of the {110}α, {200}α, {211}α, {220}α, {111}γ, {200}γ, {220}γ and {311}γ peaks were obtained and used to calculate the fγR values according to the SAE method [20]. Selected heat treated samples for electron backscatter diffraction (EBSD) analysis were ground and polished with a final vibrational polishing step using a 0.05 μm colloidal silica solution for 2 h. EBSD was performed using an EDAX-TSL detector with OIM Data Collection 5.1 software and a field emission-scanning electron microscope (FESEM) with an accelerating voltage of 20 kV, a working distance of 18 mm, and a step size of 40 nm on a hexagonal grid for each scan area of 20 × 20 μm. The results were analyzed using TSL OIM Image Analysis 7.1 software. 3. Modeling The application of a modified rule-of-mixtures accounts for load distributions between different elements in a microstructure and was initially envisioned to describe properties, such as elastic stiffness, where all elements in the analysis (i.e. constituent properties and volume fractions) were assumed constant, along with the assumption of isostrain. However, multiple analyses [21–29] have shown that the rule-ofmixtures can be more widely applied to describe more complex deformation conditions where material properties or constituent volume fractions vary with some imposed parameter, e.g. strain. Many of the current analyses have evolved from the early publication of Mileiko [21] and Garmong and Thompson [22] which incorporated strain-dependent expressions for flow stress and applied the Considère's criterion (dσ/dε = σ) [30] as an instability condition to accurately predict UE and UTS of composites. The analysis of the deformation behavior of multi-phase steels with distributed (i.e. not fibrous or layered) constituents [23–25] can be accurately captured by applying the approach incorporated in references [21,22]. Recently, applications of rule-of-mixture based models have been used to assess mechanical properties of 3G-AHSS having a variety of multi-phase microstructures which contain retained austenite that transforms to martensite during deformation [26–29]. In these analyses, multiple different functions which describe the change in constituent properties with strain (e.g. true flow stress = function of true strain) and changes in constituent volume fractions (e.g. strain-dependent transformation of austenite to martensite) have been successfully incorporated to describe overall flow equations for multi-component composites. In the present study, a model based on the rule-of mixture (Eq. (1)) [1,12] was applied for the prediction of tensile properties of multi-phase medium Mn steels, where σ is true stress and σi and Vi are true stress and volume fraction for component i, respectively. Each phase is
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Table 2 Ultimate tensile strength (UTS), strain hardening exponent (n) and strength coefficient (K) for specimens with ferrite, martensite and austenite microstructures [31–33]. Constituent
UTS (MPa)
Strain hardening exponent (n)
Strength coefficient (K)
Ref.
Ferrite Martensite Austenite
600 1750 620
0.12 0.049 0.42
873 2131 1358
Tsuchida et al. [31] Clarke [32] Poling [33]
regarded as a component distributed in a composite system. The relationship between true stress and true strain for each phase was assumed to follow the Hollomon power law (Eq. (2)), where K is the strength coefficient and n is the strain hardening exponent for each phase. σ ¼ ∑σ i Vi
ð1Þ
σ ¼ Kε n
ð2Þ
The K and n values for ferrite, martensite and austenite were obtained from a low C 0.15C–0.4Si–1.5Mn steel with a grain size of 1.5 μm [31], a water quenched 0.19C–1.6Mn–1.6Si steel [32], and an austenitic 316 L stainless steel with a chemical composition of 0.02C–1.4Mn–0.5Si– 10Ni–16Cr–2Mo and a grain size of 3.1 μm [33], respectively, and are provided in Table 2. The instability condition is obtained using Considère's criterion [30]. To predict the tensile properties of intercritically annealed medium Mn steels, estimated phase fractions (after annealing) and martensite transformation kinetics were incorporated in the composite model [1,12]. To predict the fraction and composition of each phase in the annealed specimen, a retained austenite prediction model assuming orthoequilibrium alloy partitioning at intercritical annealing temperatures was used [16]. Equilibrium fractions of selected phases (austenite, ferrite and cementite) and their chemical compositions as a function of the intercritical annealing temperature for the three alloys were calculated using Thermo-Calc® software with the TCFE 7 database. The fraction of fresh martensite transformed during quenching was then calculated by the Koistinen-Marburger (K-M) equation: f M ¼ f γ ½1 expf0:011ðM S RT Þg
ð3Þ
where fM is the fraction of martensite transformed during quenching, fγ is the intercritical austenite fraction, RT is room temperature (25 °C) and
Fig. 2. Calculated phase fractions of ferrite (α), retained austenite (γR), cementite (θ) and martensite (M) for the (a) Base, (b) Lo-Al and (c) Hi-Al steels annealed in the intercritical temperature range using the Thermo-Calc® TCFE 7 database and a retained austenite prediction model [16,17]. For comparison, experimentally measured retained austenite (γR(exp.)) values heat treated for 1 h and water quenching are also provided. Fig. 1. Relationship between MS temperature and kS value, a parameter related to the strain-induced martensite transformation kinetics in Eq. (5). The MS temperatures are calculated from austenite chemical composition and kS values are fitted using the martensite transformation kinetics during tensile deformation of the 0.1C–7.1Mn–0.1Si steel after intercritical annealing at temperatures from 575 to 650 °C for 168 h followed by water quenching [10].
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MS is the martensite start temperature calculated by the following equation [34]: MS ð CÞ ¼ 539 423C 30:4Mn 7:5Si þ 30Al
ð4Þ
The final fγR after quenching to room temperature at each intercritical annealing temperature was obtained by the subtraction of fM from fγ. The extent of strain-induced martensite transformation during tensile deformation was estimated using retained austenite chemical compositions. For a simple fitting of the martensite transformation kinetics, the Sugimoto equation (Eq. (5)) having a single term (kS) related to the austenite stability was selected [35]: f MD ¼ f γR ð1 expðkS εÞÞ
ð5Þ
where fMD is the fraction of martensite transformed during tensile deformation, ε is true strain, and kS is a parameter related to austenite stability and thus composition dependent. While there exist some studies pertaining to martensite transformation kinetics during tensile deformation in intercritically annealed medium Mn steels, retained austenite chemical compositions are rarely available [6,10,11,14]. In the present study, data for a 0.1C–7.1Mn–0.1Si steel annealed at 575, 600, 625 and 650 °C for 168 h and water quenched were used [10]. Martensite fractions during tensile deformation were measured by in situ neutron diffraction and retained austenite chemical composition was predicted using the Thermo-Calc® TCFE 2 database. Due to the limited number of chemical composition sets for the selected retained austenite conditions, the MS temperature was selected as a measure of the retained austenite stability and the relationship between the kS parameter and MS
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was determined. After determination of the MS temperature from the austenite chemical composition using Eq. (4), and fitting the kS values from experimentally measured martensite transformation kinetics as shown in Fig. 1, an exponential function was used to obtain the relationship between the MS and kS values as follows: kS ¼ 3:06 expðM S =79:7Þ 1:39:
ð6Þ
If retained austenite chemical composition is provided, this relationship provides the measure to obtain the martensite transformation kinetics during tensile deformation in combination with Eqs. (4) and (5). 4. Results and discussion Fig. 2a to c show the calculated phase fractions in the Base, Lo-Al and Hi-Al steels after intercritical annealing along with experimentally measured retained austenite (γR(exp.)) contents discussed below. The ferrite, austenite, cementite and martensite phases are denoted as α, γR, θ and M, respectively. Four critical temperatures which describe the expected phase transformation behavior are shown in each figure by vertical lines: the intercritical austenite formation temperatures (A1 and A3), the cementite dissolution temperature (ACM), and an intercritical annealing temperature where the Ms. of the intercritical austenite equals room temperature due to solute partitioning (M0) [17]. At temperatures above A3, fully austenitic microstructures exist at the annealing temperature and most of the austenite transforms to fresh martensite during quenching. Ferrite and cementite are present below the intercritical temperatures of A3 and ACM, respectively. The intercritical austenite is fully
Fig. 3. Electron backscatter diffraction (EBSD) phase maps for the Base and Hi-Al steels annealed at 700 and 725 °C for 1 h followed by water quenching: (a) Base 700 °C, (b) Base 725 °C, (c) Hi-Al 700 °C, and (d) Hi-Al 725 °C. Red and green colors denote bcc and fcc phase regions, respectively, and bcc regions include both recrystallized ferrite and fresh martensite. Arrows indicate austenite grains finely intermixed with fresh martensite (M/A constituent).
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stabilized by solute partitioning at temperatures below M0, and no martensite is present below M0. While Al additions increase the four critical temperatures, the A3 temperature shows the greatest increase from 733 °C for the Base alloy to 853 °C for the Hi-Al alloy. With Al additions, the intercritical temperature range between A1 and A3 is widened from 184 °C for the Base alloy to 282 °C for the Hi-Al alloy and the range shifts to higher temperatures. The difference between the M0 and ACM temperatures shrinks with increasing Al from 31 °C for the Base alloy to 5 °C for the Hi-Al alloy. The predicted fγR peaks shift to higher temperatures and greater values. The maximum predicted fγR values for the Base, Lo-Al and Hi-Al steels are obtained at the ACM temperature for these alloys, with values of 22.4, 25.8 and 30.8%, respectively. For comparison with the predictions, experimentally measured fγR values for the three steels at temperatures ranging from 550 to 725 °C for 1 h annealing are also provided in each figure. Experimentally measured maximum fγR values for the Base, Lo-Al and Hi-Al steels are 14.1, 17.1 and 22.3%, and corresponding temperatures for the maximum retained austenite contents are 675, 675 and 700 °C, respectively. In comparison to predicted values, the experimentally measured maximum fγR values are lower and annealing temperatures are higher, although the influence of Al displays the predicted trend. With an increase in Al content, both the predicted and experimental results exhibit increased peak fγR values and peak annealing temperatures. The amount of martensite in the Base and Hi-Al steels was examined in the samples annealed at relatively high temperatures (700 and 725 °C) using EBSD phase maps as shown in Fig. 3a to d. The red and green colors denote bcc and fcc phase regions, respectively. While the bcc indexed regions indicate either recrystallized ferrite or fresh martensite, the two microstructures have different morphologies. Recrystallized ferrite has a polygonal grain structure, whereas fresh martensite is finely distributed in a retained austenite matrix. The Base steel samples annealed at 700 °C (Fig. 3a) and 725 °C (Fig. 3b) exhibit a mixed microstructure composed of polygonal bcc grains (recrystallized ferrite), fine austenite grains associated with fine bcc regions (martensite/austenite (M/A) constituents, white arrows), and a small fraction of large equiaxed austenite grains. While the Hi-Al steel annealed at 700 °C (Fig. 3c) seems to predominantly exhibit large equiaxed austenite and ferrite grains, some of the austenite grains in the Hi-Al steel annealed at 725 °C (Fig. 3d) appear associated with the M/A constituents. These M/A constituents support the predictions that fresh martensite forms during quenching, and that more fresh martensite would be present in the Base steel than in the Hi-Al steel.
To confirm the interpretation related to fresh martensite formation, Fig. 4 shows bcc image quality (IQ) distributions generated for the Base and Hi-Al alloys from EBSD data. Each plot shows the presence of two peaks interpreted to reflect the two possible bcc microstructures, i.e. recrystallized ferrite (high IQ peak) and fresh martensite (low IQ peak).
Fig. 4. Electron backscatter diffraction (EBSD) image quality (IQ) distributions of bcc phase in the Base and Hi-Al steels annealed at 700 and 725 °C for 1 h followed by water quenching. The intensities of the peaks at lower IQ and higher IQ reflect the fractions of martensite and ferrite, respectively.
Fig. 5. Equilibrium (a) C, (b) Mn and (c) Al contents in the intercritical austenite for the Base, Lo-Al and Hi-Al steels as a function of annealing temperature predicted by the Thermo-Calc® TCFE 7 database. The critical temperatures (A1, M0, ACM and A3) are shown (only) for the Hi-Al alloy.
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Due to its tetragonality and high dislocation density, the martensite exhibits relatively low IQ values compared to ferrite. For each alloy, the relative intensity of the lower IQ peak decreases at lower annealing temperature. In particular, the lower IQ peak almost disappears in the Hi-Al 700 °C sample, implying less martensite formation during quenching. In both steels, martensite fractions (i.e., magnitude proportional to peak height) increase with annealing temperature, consistent with the predicted results shown in Fig. 2. The retained austenite stabilities for the three alloys with different annealing temperatures were estimated from the retained austenite chemical compositions. Calculated equilibrium amounts of C, Mn and Al in the intercritically annealed austenite for the three alloys are plotted against the intercritical annealing temperature in Fig. 5a to c, assuming full partitioning. The four critical temperatures provided in each figure are for the Hi-Al alloy. With decreasing annealing temperature, Fig. 5a shows that the predicted austenite carbon content increased below A3, showed a maximum at ACM and decreased due to cementite precipitation below ACM. Equilibrium partitioning of Mn (Fig. 5b) is clearly greater at lower annealing temperatures where less austenite and more ferrite are present. The austenite Al content correspondingly decreases at lower intercritical annealing temperatures (Fig. 5c). To estimate the mechanical stability of the retained austenite for the three alloys, the MS temperatures and kS values were calculated using Eqs. (4) and (6) based on the chemical compositions predicted at each annealing temperature, and the results are shown in Fig. 6. Over the entire intercritical annealing temperature range, the MS temperatures of the three alloys gradually decrease with decreasing annealing temperature (Fig. 6a). Increased austenite stabilization originates from the
Fig. 6. Predicted (a) MS temperatures and (b) kS values of the retained austenite in the Base, Lo-Al and Hi-Al steels after intercritical annealing.
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enrichment of C and Mn and depletion of Al. The austenite stability exhibits increased temperature dependence in the Hi-Al alloy. Based on the MS temperature in Fig. 6a, Fig. 6b shows the calculated kS values using Eq. (4). Similar to the MS temperature, the kS values decrease with lower annealing temperature and Al additions amplified the reduction. Lower kS values imply less martensite transformation with strain. Fig. 7a to c show the measured room temperature tensile flow curves for the Base, Lo-Al and Hi-Al steels annealed at temperatures ranging from 550 to 725 °C for 1 h followed by water quenching. In each alloy, tensile deformation characteristics are strongly influenced by the annealing temperature. For example, the total elongation in the Base alloy annealed at 550 °C is 4.5%, and greatly increased up to 38.8% for the sample annealed at 650 °C. In addition, the maximum UTS value of 1603 MPa obtained at 725 °C is twice as large as the
Fig. 7. Room temperature engineering stress-strain curves of the (a) Base, (b) Lo-Al and (c) Hi-Al steels annealed at temperatures ranging from 550 to 725 °C for 1 h followed by water quenching.
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minimum value of 730 MPa obtained at 625 °C. As expected, samples annealed at temperatures between 650 and 700 °C having higher fγR values exhibited relatively larger elongations, and greater UTS was observed in the samples annealed above 700 °C having higher martensite fractions. The slopes of the tensile flow curves tend to increase with annealing temperature, implying increased strain hardening, reflecting increased transformation of austenite to martensite with strain. This behavior is consistent with predictions in Fig. 6b, where kS values increased with temperature, indicating more rapid martensite transformation during strain for microstructures processed at higher annealing temperatures, providing increased strain hardening. Similar evolutions of tensile flow curves with annealing temperatures are observed in both the Lo-Al (Fig. 7b) and the Hi-Al alloys (Fig. 7c). Experimental and predicted UTS and UE values as a function of annealing temperature for the Base, Lo-Al and Hi-Al steels are provided in Fig. 8. Fig. 8a shows that with an increase in annealing temperature, experimentally measured UTS values for the three alloys initially decrease, exhibit a minimum at temperatures near 650 °C, and dramatically increase at higher temperatures. For all tested conditions, Al additions decrease the predicted UTS values and the effects of Al additions are greater at higher annealing temperatures. Fig. 8b shows the corresponding UTS values predicted for the Base, Lo-Al and Hi-Al steels. The predictions exhibit similar variations with annealing temperatures as the experimental data. Predicted temperatures corresponding to the lowest UTS values in the Base, Lo-Al and Hi-Al steels are 610, 632 and 666 °C, respectively, which are equal or slightly less than the M0 temperature values (614, 632 and 666 °C for the Base, Lo-Al and Hi-Al steels, respectively), below which martensite formation is not anticipated during quenching from the intercritical heat treating temperature.
In Fig. 8c, experimentally measured UE values for the Base, Lo-Al and Hi-Al steels exhibit a peak as a function of annealing temperature. Al additions decrease the maximum UE values from 35.2 to 25.6% and shift the peaks to higher temperatures. Predicted UE values for the Base, Lo-Al and Hi-Al steels shown in Fig. 8d exhibit the maximum UE at 620, 640 and 670 °C, respectively. These temperatures are slightly higher than the temperatures for the lowest UTS. Overall, the predicted UE values with annealing temperature and Al additions show reasonable agreement with the experimental results. However, Al additions increase the maximum predicted UE value from 20.3, 22.1 to 25.3%, in contrast to the experimental results. One of the possible reasons for this discrepancy is the composition dependent phase tensile properties. While not considered for the predictions in the present study, tensile properties of each phase are influenced by their chemical compositions. Fig. 9a shows the relationship between UTS×UE values and annealing temperature for the Base, Lo-Al and Hi-Al steels obtained from predictions and experiments. Such products are often used to compare the overall strength-ductility balance for different microstructures. Overall, experimental UTS×UE values show higher maximum values and larger variations with annealing temperature compared to the predicted results. The experimental UTS×UE values for the Base, Lo-Al and Hi-Al alloys exhibit a maximum at 675, 675 and 700 °C, respectively, and predicted maximum values are obtained at 645, 654 and 671 °C (corresponding to the ACM for each alloy), respectively. For both experimental and predicted results, the maximum UTS×UE values are obtained at temperatures having the maximum fγR values. The relationships between the UTS×UE and fγR values for the three alloys are shown in Fig. 9b. While both predicted and experimental results show a strong relationship between the UTS×UE and fγR values, the effect of the fγR
Fig. 8. Tensile properties of the Base, Lo-Al and Hi-Al steels: (a) experimentally measured ultimate tensile strength (UTS, Exp.), (b) predicted ultimate tensile strength (UTS, Pre.), (c) experimentally measured uniform elongation (UE, Exp.), and (d) predicted uniform elongation (UE, Pre.).
S. Kang et al. / Materials and Design 97 (2016) 138–146
on the UTS×UE is somewhat underestimated in the predicted results. Also, the results in Fig. 9b illustrate the importance of retained austenite on the tensile properties of intercritically annealed medium Mn steels. Medium Mn steels after intercritical annealing have diverse combinations of various phases and austenite stability, resulting in large variations of tensile deformation behavior. The model introduced in the present study provides an important means of predicting tensile properties of medium Mn steels using predicted phase fractions and TRIP behavior using only thermodynamic inputs, without requiring any direct experimental observations. Through this modeling method, it is possible to estimate the complex effects of alloy composition and annealing temperature on the tensile deformation behavior of medium Mn steels. While the accuracy of the predicted results can be improved, the models should be useful in the design of alloy compositions and heat treatment processes for improved tensile properties of medium Mn steels. 5. Conclusions The tensile properties of intercritically annealed Base, Lo-Al and Hi-Al steels were predicted and the ultimate tensile strengths (UTS) and uniform elongations (UE) were compared to experimental data for three 0.19C, 4.5Mn steels with different Al additions. Intercritically annealed phase fractions and retained austenite chemical compositions were predicted using a retained austenite prediction model assuming ortho-equilibrium solute partitioning. The martensite transformation behavior during tensile deformation was estimated from the predicted retained austenite chemical composition. Tensile properties, i.e. UTS and UE of intercritically annealed medium Mn steels were predicted
Fig. 9. Ultimate tensile strength (UTS)-uniform elongation (UE) products for the Base, LoAl and Hi-Al steels obtained from experiments (Exp.) and predictions (Pre.) as a function of (a) intercritical annealing temperature and (b) retained austenite fraction (fγR).
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by incorporating the predicted phase fractions and martensite transformation kinetics in a simplified composite model for multi-component systems. Predictions show minimum UTS values near the M0 temperature, maximum UE values at temperatures between the M0 and ACM, and maximum UTS×UE values at ACM where the greatest retained austenite fractions were predicted. Overall, the evolution of predicted tensile properties with annealing temperature and Al additions exhibited reasonable agreement with the experimental data, reflecting the importance of retained austenite, and austenite stability. Acknowledgement The authors gratefully acknowledge the support of the sponsors of the Advanced Steel Processing and Products Research Center, an industry/university cooperative research center at the Colorado School of Mines. References [1] D.K. Matlock, J.G. Speer, E. De Moor, P.J. Gibbs, Recent developments in advanced high strength sheet steels for automotive applications: an overview, Jestech 15 (2012) 1–12. [2] I. Tamura, Deformation-induced martensitic transformation and transformationinduced plasticity in steels, Met. Sci. 16 (1982) 245–253. [3] R.L. Miller, Ultrafine-grained microstructures and mechanical properties of alloy steels, Metall. Trans. A 3 (1972) 905–912. [4] Y.-K. Lee, J. Han, Current opinion in medium manganese steel, Mater. Sci. Technol. 31 (2015) 843–856. [5] S. Lee, B.C. De Cooman, On the selection of the optimal intercritical annealing temperature for medium Mn TRIP steel, Metall. Mater. Trans. A 44 (2013) 5018–5024. [6] D.-W. Suh, S.-J. Park, T.-H. Lee, C.-S. Oh, S.-J. Kim, Influence of Al on the microstructural evolution and mechanical behavior of low-carbon, manganese transformationinduced-plasticity steel, Metall. Mater. Trans. A 41 (2010) 397–408. [7] T. Furukawa, Dependence of strength-ductility characteristics on thermal history in low carbon, 5 wt-%Mn steels, Mater. Sci. Technol. 5 (1989) 465–470. [8] H. Huang, O. Matsumura, T. Furukawa, Retained austenite in low carbon, manganese steel after intercritical heat treatment, Mater. Sci. Technol. 10 (1994) 621–626. [9] D.E. Kim, Y.K. Park, O.Y. Lee, K.G. Jin, S.J. Kim, Formation of retained austenite and mechanical properties of 4–8%Mn hot rolled TRIP steels, Korean J. Mater. Res. 15 (2005) 115–120. [10] P.J. Gibbs, E. De Moor, M.J. Merwin, B. Clausen, J.G. Speer, D.K. Matlock, Austenite stability effects on tensile behavior of manganese-enriched-austenite transformationinduced plasticity steel, Metall. Mater. Trans. A 42 (2011) 3691–3702. [11] D.W. Suh, J.H. Ryu, M.S. Joo, H.S. Yang, K. Lee, H.K.D.H. Bhadeshia, Medium-alloy manganese-rich transformation-induced plasticity steels, Metall. Mater. Trans. A 44 (2013) 286–293. [12] D.K. Matlock, J.G. Speer, Design considerations for the next generation of advanced high strength steels, in: H.C. Lee (Ed.), Proceedings of the 3rd International Conference on Advanced Structural Steels, Kor. Inst. Metal. Mater., Seoul, South Korea 2006, pp. 774–781. [13] R. Rana, P.J. Gibbs, E. De Moor, J.G. Speer, D.K. Matlock, A composite modeling analysis of the deformation behavior of medium manganese steels, Steel Res. Int. 85 (2015) 1–12. [14] S. Lee, Y. Estrin, B.C. De Cooman, Constitutive modeling of the mechanical properties of V-added medium manganese TRIP steel, Metall. Mater. Trans. A 44 (2013) 3136–3146. [15] Y. Estrin, H. Mecking, A unified phenomenological description of work hardening and creep based on one-parameter models, Acta Mater. 32 (1984) 57–70. [16] E. De Moor, D.K. Matlock, J.G. Speer, M.J. Merwin, Austenite stabilization through manganese enrichment, Scr. Mater. 64 (2011) 185–188. [17] S. Kang, E. De Moor, J.G. Speer, Retained austenite stabilization through solute partitioning during intercritical annealing in C-, Mn-, Al-, Si-, and Cr-alloyed steels, Metall. Mater. Trans. A 46 (2015) 1005–1011. [18] S. Kang, E. De Moor, J.G. Speer, D. Krizan, in: M. Militzer, G. Botton, L.-Q. Chen, J. Howe, C. Sinclair, H. Zurob (Eds.),Aluminum alloyed medium manganese steels, Proceedings of the International Conference on Solid-Solid Phase Transformations in Inorganic Materials 2015 2015, pp. 1145–1151 (Whistler, BC, Canada). [19] ASTM E8/E8M-11, Standard Test Methods for Tension Testing of Metallic Materials, West Conshohocken, PA, ASTM International, 2011. [20] C.F. Jatczak, Retained austenite and its measurement by X-ray diffraction, SAE Technical Paper No. 800426, 1980. [21] S.T. Mileiko, The tensile strength and ductility of continuous fibre composites, J. Mater. Sci. 4 (1969) 974–977. [22] G. Garmong, R.B. Thompson, A theory for the mechanical properties of metal-matrix composites at ultimate loading, Metall. Trans. 4 (1973) 863–873. [23] R.G. Davies, The deformation behavior of a vanadium-strengthened dual-phase steel, Metall. Trans. A 9A (1978) 41–52. [24] G.R. Speich, R.L. Miller, Mechanical properties of ferrite–martensite steels, in: R.A. Kot, J.W. Morris (Eds.), Structure and Properties of Dual-phase Steels, TMS, Warrendale, PA 1979, pp. 145–182.
146
S. Kang et al. / Materials and Design 97 (2016) 138–146
[25] J. Becker, E. Hornbogen, F. Wendl, Analysis of mechanical properties of a low-alloyed mo-steel with different phosphorus additions, in: R.A. Kot, B.L. Bramfitt (Eds.), Fundamentals of Dual-Phase Steels, TMS, Warrendale PA 1981, pp. 383–398. [26] S. Lee, Y. Estrin, B.C. De Cooman, Constitutive modeling of the mechanical properties of V-added medium manganese TRIP steel, Metall. Mater. Trans. A 44A (2013) 3136–3146. [27] H.N. Han, C.S. Oh, G. Kim, O. Kwon, Design method for TRIP-aided multiphase steel based on microstructure-based modelling for transformation induced plasticity and mechanically induced martensitic transformation, Mater. Sci. Eng. A 499 (2009) 462–468. [28] O. Bouaziz, J.D. Embury, Microstructural design for advanced structural steels, Mater. Sci. Forum 539-543 (2007) 42–50. [29] K.S. Choi, A. Soulami, W.N. Liu, X. Sun, M.A. Khaleel, Influence of various material design parameters on deformation behaviors of TRIP steels, Comput. Mater. Sci. 50 (2010) 720–730.
[30] G.E. Dieter, Mechanical Metallurgy 3rd Edition, McGraw-Hill, London, 1986. [31] N. Tsuchida, H. Masuda, Y. Harada, K. Fukaura, Y. Tomota, K. Nagai, Effect of ferrite grain size on tensile deformation behavior of a ferrite-cementite low carbon steel, Mater. Sci. Eng. A 488 (2008) 446–452. [32] A. Clarke, Carbon Partitioning Into Austenite From Martensite in a Silicon-containing High Strength Sheet Steel(PhD Thesis) Colorado School of Mines, 2006. [33] W.A. Poling, Grain Size Effects in Micro-tensile Testing of Austenitic Stainless Steel(MS Thesis) Colorado School of Mines, 2012. [34] K.W. Andrews, Empirical formulae for the calculation of some transformation temperatures, J. Iron Steel Inst. 203 (1965) 721–727. [35] K. Sugimoto, M. Kobayashi, S. Hashimoto, Ductility and strain-induced transformation in a high-strength transformation-induced plasticity-aided dual-phase steel, Metall. Mater. Trans. A 23 (1992) 3085–3091.
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Prediction of tensile properties of intercritically annealed Al-containing 0.19C–4.5Mn (wt%) TRIP steels Singon Kang a, John G. Speer a, Daniel Krizan b, David K. Matlock a, Emmanuel De Moor a,⁎ a b
Advanced Steel Processing and Products Research Center, Colorado School of Mines, Golden, CO 80401, USA Research and Development Department Business Unit Coil, Voestalpine Steel Division GmbH, Voestalpine-Str. 3, Linz A-4031, Austria
a r t i c l e
i n f o
Article history: Received 3 November 2015 Received in revised form 8 February 2016 Accepted 15 February 2016 Available online 17 February 2016 Keywords: Transformation-induced plasticity Retained austenite Composite model Intercritical annealing Alloy partitioning
a b s t r a c t Tensile properties were predicted for intercritically annealed 4.5 wt% Mn steels with varying Al additions up to 1.3 wt%. Intercritically annealed phase fractions and retained austenite chemical compositions were obtained using a retained austenite prediction model assuming ortho-equilibrium alloy partitioning. The martensite transformation kinetics during tensile deformation were also estimated from the predicted retained austenite chemical compositions. Predicted phase fractions and martensite transformation kinetics were incorporated in a composite model for multi-component systems. Predicted tensile strength and uniform elongation values were compared to the experimental data obtained from the steels annealed at temperatures ranging from 550 to 725 °C for 1 h followed by water quenching. © 2016 Elsevier Ltd. All rights reserved.
1. Introduction Recent developments in “third generation” advanced high strength steels (3G-AHSS) for automotive applications aim to achieve high strength and ductility to meet the needs for fuel efficiency and passenger safety. Phase transformation from austenite to martensite during tensile deformation, i.e. transformation-induced plasticity (TRIP), is one of the prominent strategies to obtain desired tensile properties for 3G-AHSS [1]. Through the strain-induced or stress-assisted transformation of austenite to the harder martensite phase, exceptional strain hardening occurs during tensile deformation, improving formability of these high strength steels [2]. Since the work by Miller [3] on a 0.11C–5.7Mn steel, medium Mn steels have been extensively investigated [4–14]. Heat treatment of medium Mn steels in the intercritical temperature range enables partitioning of alloying elements between ferrite and austenite. Austenite enrichment by stabilizing elements such as Mn and C restricts martensite formation during quenching and increases retained austenite fractions (fγR). To obtain reasonable amounts of fγR by intercritical annealing and understand the effects of retained austenite on tensile properties, studies have been performed using medium Mn TRIP steels [7–9]. Huang et al. [8] annealed a 0.12C–5.1Mn steel using a variety of temperatures, holding times and cooling conditions. For 3 h annealing followed by a water quench, a variation in annealing temperature from 700 to ⁎ Corresponding author. E-mail address: [email protected] (E. De Moor).
http://dx.doi.org/10.1016/j.matdes.2016.02.058 0264-1275/© 2016 Elsevier Ltd. All rights reserved.
650 °C increased the fγR from about 10 to 30%, and improved the product of ultimate tensile strength (UTS) and uniform elongation (UE) values from 10,000 to 25,000 MPa·%. Kim et al. [9] reported that the UTS×UE values for 0.1 C steels containing 4, 6 or 8 wt% Mn annealed in the intercritical temperature range, increased with an increase in fγR. This confirms that tensile properties of medium Mn TRIP steels may be enhanced by increased fγR. The importance of austenite fractions has been reported by others as well, along with austenite stability [4,5,10,11]. The effect of martensite transformation behavior during tensile deformation on the tensile properties of intercritically annealed medium Mn steels has been examined e.g. [10,11]. Gibbs et al. [10] performed intercritical annealing of a 0.1C–7.1Mn–0.1Si steel in a temperature range of 575 to 675 °C for 168 h followed by water quenching. Although less austenite was stabilized by annealing at 575 °C (26.3%) compared to 650 °C (43.5%), the 650 °C sample exhibited reduced elongation values compared to the 575 °C sample. This behavior was ascribed to the rapid transformation kinetics in the 650 °C sample as more than 80% of the retained austenite transformed to martensite within a true strain range of 0.05. A similar reduction in austenite stability with increasing annealing temperature was observed by Suh et al. [11] in a 0.11C– 4.5Mn–0.5Si–2.2Al steel. In addition to the amount and stability of austenite, the fractions of other phases such as ferrite, cementite and martensite are also greatly influenced by the intercritical annealing conditions [10], and correspondingly the resulting tensile properties depend on the ferrite, cementite and martensite fractions in annealed medium Mn steels. While there are multiple approaches currently being considered to predict the properties of multi-component systems which evolve with
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imposed macroscopic strain to account for strain and stress gradients between individual constituents, a modified rule-of-mixture approach is frequently used for predictions of tensile properties and it provides an insight into design requirements for new steels with complex microstructures. Matlock and Speer [1,12] used the modified rule-of-mixture approach to propose a composite model for predictions of tensile properties incorporating martensite transformation kinetics during tensile deformation. The tensile flow curves for each phase were assumed to follow Hollomon's power law, and the strain dependent phase fractions were applied to consider the effect of martensite transformation during the tensile deformation. Predicted variations in tensile properties of multi-phase steels were provided for different combinations of phase fractions and martensite transformation kinetics. Rana et al. [13] applied the composite model to an experimental alloy having a chemical composition of 0.1C–7.1Mn–0.1Si annealed at 600, 625 and 650 °C. Their phase fractions and martensite transformation kinetics were experimentally measured [10] and model predictions were shown to directly predict observed strengths and ductilities. Lee et al. [14] predicted the tensile properties of a 0.08C–6Mn– 2.0Al–1.5Si–0.08V steel annealed at 740 °C for 3 min using a modified Estrin and Mecking constitutive model [15]. Respective constitutive models for each phase such as ferrite, austenite, and martensite were constructed using diverse experimental parameters including dislocation density, grain size, and martensite transformation kinetics. These constitutive models were combined by considering strain partitioning between the phases based on iso-work theory. These predictions successfully describe the need for significant amounts of retained austenite with controlled stability against deformation for new 3G-AHSS. However, the application of these models was limited for sample conditions with existing experimental results such as initial phase fraction and martensite transformation kinetics. In the present study, a method is proposed to predict the tensile properties of intercritically annealed multi-phase steels, which does not require pre-existing experimental data. A composite model proposed by Matlock and Speer [1,12] which incorporates predicted initial phase fraction and martensite transformation kinetics is used to obtain the tensile properties. Phase fractions and compositions for annealed samples are predicted using a model assuming ortho-equilibrium alloy partitioning during intercritical annealing [16,17]. Martensite transformation kinetics during tensile deformation is estimated based on the retained austenite chemical composition. This method is applied to three medium Mn steels containing different levels of Al with varying intercritical annealing temperature, which have different combinations of various phases and a wide range of fγR. For verification of the proposed model, predicted results are compared to experimental tensile properties obtained from equivalent alloy and heat treatment conditions. The evolution of the tensile properties with annealing temperature in both prediction and experiments are compared and discussed. 2. Experimental procedure Three 0.19C–4.5Mn (wt%) medium Mn steels containing 0, 0.5 or 1.3 wt% Al, referred to as Base, Lo-Al and Hi-Al, having the chemical compositions provided in Table 1, were designed and laboratory cast as 90 kg ingots [18]. Sample blocks having a thickness of 130 mm were reheated to 1180 °C for 2 h, hot rolled in 12 passes to a thickness of 4 mm with a finishing temperature of 900 °C, and coiled at 500 °C. These hot rolled plates had fully martensitic microstructures. In order Table 1 Chemical compositions of the Base, Lo-Al and Hi-Al alloys. wt%
C
Mn
Al
Base Lo-Al Hi-Al
0.190 0.189 0.195
4.39 4.45 4.52
0.01 0.48 1.30
139
to facilitate cold rolling without crack formation, the plates were softened by a prior batch annealing at 550 °C for 16 h where the microstructure was primarily ferrite and thus alloy partitioning to austenite would be minimized. These softened plates were cold rolled in 5 passes to a final thickness of 1 mm. Tensile specimens were machined along the rolling direction with a gauge length of 50.4 mm according to the ASTM-E8 standard [19]. The specimens were isothermally heat treated at temperatures ranging from 550 to 725 °C for 1 h in a box furnace and water quenched. To avoid sample oxidation, the samples were sealed in stainless steel bags with Ti particles to getter oxygen. Tensile tests for the heat treated samples were performed at a constant strain rate of 6.25 × 10−4 s−1 at room temperature. Heat treated samples for X-ray diffraction (XRD) analysis were ground with 1200 grit SiC paper and chemically polished with a 1:50:50 mixture of 48% hydrofluoric acid, 30% hydrogen peroxide, and de-ionized water at room temperature for 180 s. A Philips X-pert diffractometer with a copper tube was operated at 45 kV and 40 mA and a 2θ scan range from 40 to 110° was employed. After stripping the Kα2diffracted peaks through X-Pert software, the intensities of the {110}α, {200}α, {211}α, {220}α, {111}γ, {200}γ, {220}γ and {311}γ peaks were obtained and used to calculate the fγR values according to the SAE method [20]. Selected heat treated samples for electron backscatter diffraction (EBSD) analysis were ground and polished with a final vibrational polishing step using a 0.05 μm colloidal silica solution for 2 h. EBSD was performed using an EDAX-TSL detector with OIM Data Collection 5.1 software and a field emission-scanning electron microscope (FESEM) with an accelerating voltage of 20 kV, a working distance of 18 mm, and a step size of 40 nm on a hexagonal grid for each scan area of 20 × 20 μm. The results were analyzed using TSL OIM Image Analysis 7.1 software. 3. Modeling The application of a modified rule-of-mixtures accounts for load distributions between different elements in a microstructure and was initially envisioned to describe properties, such as elastic stiffness, where all elements in the analysis (i.e. constituent properties and volume fractions) were assumed constant, along with the assumption of isostrain. However, multiple analyses [21–29] have shown that the rule-ofmixtures can be more widely applied to describe more complex deformation conditions where material properties or constituent volume fractions vary with some imposed parameter, e.g. strain. Many of the current analyses have evolved from the early publication of Mileiko [21] and Garmong and Thompson [22] which incorporated strain-dependent expressions for flow stress and applied the Considère's criterion (dσ/dε = σ) [30] as an instability condition to accurately predict UE and UTS of composites. The analysis of the deformation behavior of multi-phase steels with distributed (i.e. not fibrous or layered) constituents [23–25] can be accurately captured by applying the approach incorporated in references [21,22]. Recently, applications of rule-of-mixture based models have been used to assess mechanical properties of 3G-AHSS having a variety of multi-phase microstructures which contain retained austenite that transforms to martensite during deformation [26–29]. In these analyses, multiple different functions which describe the change in constituent properties with strain (e.g. true flow stress = function of true strain) and changes in constituent volume fractions (e.g. strain-dependent transformation of austenite to martensite) have been successfully incorporated to describe overall flow equations for multi-component composites. In the present study, a model based on the rule-of mixture (Eq. (1)) [1,12] was applied for the prediction of tensile properties of multi-phase medium Mn steels, where σ is true stress and σi and Vi are true stress and volume fraction for component i, respectively. Each phase is
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Table 2 Ultimate tensile strength (UTS), strain hardening exponent (n) and strength coefficient (K) for specimens with ferrite, martensite and austenite microstructures [31–33]. Constituent
UTS (MPa)
Strain hardening exponent (n)
Strength coefficient (K)
Ref.
Ferrite Martensite Austenite
600 1750 620
0.12 0.049 0.42
873 2131 1358
Tsuchida et al. [31] Clarke [32] Poling [33]
regarded as a component distributed in a composite system. The relationship between true stress and true strain for each phase was assumed to follow the Hollomon power law (Eq. (2)), where K is the strength coefficient and n is the strain hardening exponent for each phase. σ ¼ ∑σ i Vi
ð1Þ
σ ¼ Kε n
ð2Þ
The K and n values for ferrite, martensite and austenite were obtained from a low C 0.15C–0.4Si–1.5Mn steel with a grain size of 1.5 μm [31], a water quenched 0.19C–1.6Mn–1.6Si steel [32], and an austenitic 316 L stainless steel with a chemical composition of 0.02C–1.4Mn–0.5Si– 10Ni–16Cr–2Mo and a grain size of 3.1 μm [33], respectively, and are provided in Table 2. The instability condition is obtained using Considère's criterion [30]. To predict the tensile properties of intercritically annealed medium Mn steels, estimated phase fractions (after annealing) and martensite transformation kinetics were incorporated in the composite model [1,12]. To predict the fraction and composition of each phase in the annealed specimen, a retained austenite prediction model assuming orthoequilibrium alloy partitioning at intercritical annealing temperatures was used [16]. Equilibrium fractions of selected phases (austenite, ferrite and cementite) and their chemical compositions as a function of the intercritical annealing temperature for the three alloys were calculated using Thermo-Calc® software with the TCFE 7 database. The fraction of fresh martensite transformed during quenching was then calculated by the Koistinen-Marburger (K-M) equation: f M ¼ f γ ½1 expf0:011ðM S RT Þg
ð3Þ
where fM is the fraction of martensite transformed during quenching, fγ is the intercritical austenite fraction, RT is room temperature (25 °C) and
Fig. 2. Calculated phase fractions of ferrite (α), retained austenite (γR), cementite (θ) and martensite (M) for the (a) Base, (b) Lo-Al and (c) Hi-Al steels annealed in the intercritical temperature range using the Thermo-Calc® TCFE 7 database and a retained austenite prediction model [16,17]. For comparison, experimentally measured retained austenite (γR(exp.)) values heat treated for 1 h and water quenching are also provided. Fig. 1. Relationship between MS temperature and kS value, a parameter related to the strain-induced martensite transformation kinetics in Eq. (5). The MS temperatures are calculated from austenite chemical composition and kS values are fitted using the martensite transformation kinetics during tensile deformation of the 0.1C–7.1Mn–0.1Si steel after intercritical annealing at temperatures from 575 to 650 °C for 168 h followed by water quenching [10].
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MS is the martensite start temperature calculated by the following equation [34]: MS ð CÞ ¼ 539 423C 30:4Mn 7:5Si þ 30Al
ð4Þ
The final fγR after quenching to room temperature at each intercritical annealing temperature was obtained by the subtraction of fM from fγ. The extent of strain-induced martensite transformation during tensile deformation was estimated using retained austenite chemical compositions. For a simple fitting of the martensite transformation kinetics, the Sugimoto equation (Eq. (5)) having a single term (kS) related to the austenite stability was selected [35]: f MD ¼ f γR ð1 expðkS εÞÞ
ð5Þ
where fMD is the fraction of martensite transformed during tensile deformation, ε is true strain, and kS is a parameter related to austenite stability and thus composition dependent. While there exist some studies pertaining to martensite transformation kinetics during tensile deformation in intercritically annealed medium Mn steels, retained austenite chemical compositions are rarely available [6,10,11,14]. In the present study, data for a 0.1C–7.1Mn–0.1Si steel annealed at 575, 600, 625 and 650 °C for 168 h and water quenched were used [10]. Martensite fractions during tensile deformation were measured by in situ neutron diffraction and retained austenite chemical composition was predicted using the Thermo-Calc® TCFE 2 database. Due to the limited number of chemical composition sets for the selected retained austenite conditions, the MS temperature was selected as a measure of the retained austenite stability and the relationship between the kS parameter and MS
141
was determined. After determination of the MS temperature from the austenite chemical composition using Eq. (4), and fitting the kS values from experimentally measured martensite transformation kinetics as shown in Fig. 1, an exponential function was used to obtain the relationship between the MS and kS values as follows: kS ¼ 3:06 expðM S =79:7Þ 1:39:
ð6Þ
If retained austenite chemical composition is provided, this relationship provides the measure to obtain the martensite transformation kinetics during tensile deformation in combination with Eqs. (4) and (5). 4. Results and discussion Fig. 2a to c show the calculated phase fractions in the Base, Lo-Al and Hi-Al steels after intercritical annealing along with experimentally measured retained austenite (γR(exp.)) contents discussed below. The ferrite, austenite, cementite and martensite phases are denoted as α, γR, θ and M, respectively. Four critical temperatures which describe the expected phase transformation behavior are shown in each figure by vertical lines: the intercritical austenite formation temperatures (A1 and A3), the cementite dissolution temperature (ACM), and an intercritical annealing temperature where the Ms. of the intercritical austenite equals room temperature due to solute partitioning (M0) [17]. At temperatures above A3, fully austenitic microstructures exist at the annealing temperature and most of the austenite transforms to fresh martensite during quenching. Ferrite and cementite are present below the intercritical temperatures of A3 and ACM, respectively. The intercritical austenite is fully
Fig. 3. Electron backscatter diffraction (EBSD) phase maps for the Base and Hi-Al steels annealed at 700 and 725 °C for 1 h followed by water quenching: (a) Base 700 °C, (b) Base 725 °C, (c) Hi-Al 700 °C, and (d) Hi-Al 725 °C. Red and green colors denote bcc and fcc phase regions, respectively, and bcc regions include both recrystallized ferrite and fresh martensite. Arrows indicate austenite grains finely intermixed with fresh martensite (M/A constituent).
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stabilized by solute partitioning at temperatures below M0, and no martensite is present below M0. While Al additions increase the four critical temperatures, the A3 temperature shows the greatest increase from 733 °C for the Base alloy to 853 °C for the Hi-Al alloy. With Al additions, the intercritical temperature range between A1 and A3 is widened from 184 °C for the Base alloy to 282 °C for the Hi-Al alloy and the range shifts to higher temperatures. The difference between the M0 and ACM temperatures shrinks with increasing Al from 31 °C for the Base alloy to 5 °C for the Hi-Al alloy. The predicted fγR peaks shift to higher temperatures and greater values. The maximum predicted fγR values for the Base, Lo-Al and Hi-Al steels are obtained at the ACM temperature for these alloys, with values of 22.4, 25.8 and 30.8%, respectively. For comparison with the predictions, experimentally measured fγR values for the three steels at temperatures ranging from 550 to 725 °C for 1 h annealing are also provided in each figure. Experimentally measured maximum fγR values for the Base, Lo-Al and Hi-Al steels are 14.1, 17.1 and 22.3%, and corresponding temperatures for the maximum retained austenite contents are 675, 675 and 700 °C, respectively. In comparison to predicted values, the experimentally measured maximum fγR values are lower and annealing temperatures are higher, although the influence of Al displays the predicted trend. With an increase in Al content, both the predicted and experimental results exhibit increased peak fγR values and peak annealing temperatures. The amount of martensite in the Base and Hi-Al steels was examined in the samples annealed at relatively high temperatures (700 and 725 °C) using EBSD phase maps as shown in Fig. 3a to d. The red and green colors denote bcc and fcc phase regions, respectively. While the bcc indexed regions indicate either recrystallized ferrite or fresh martensite, the two microstructures have different morphologies. Recrystallized ferrite has a polygonal grain structure, whereas fresh martensite is finely distributed in a retained austenite matrix. The Base steel samples annealed at 700 °C (Fig. 3a) and 725 °C (Fig. 3b) exhibit a mixed microstructure composed of polygonal bcc grains (recrystallized ferrite), fine austenite grains associated with fine bcc regions (martensite/austenite (M/A) constituents, white arrows), and a small fraction of large equiaxed austenite grains. While the Hi-Al steel annealed at 700 °C (Fig. 3c) seems to predominantly exhibit large equiaxed austenite and ferrite grains, some of the austenite grains in the Hi-Al steel annealed at 725 °C (Fig. 3d) appear associated with the M/A constituents. These M/A constituents support the predictions that fresh martensite forms during quenching, and that more fresh martensite would be present in the Base steel than in the Hi-Al steel.
To confirm the interpretation related to fresh martensite formation, Fig. 4 shows bcc image quality (IQ) distributions generated for the Base and Hi-Al alloys from EBSD data. Each plot shows the presence of two peaks interpreted to reflect the two possible bcc microstructures, i.e. recrystallized ferrite (high IQ peak) and fresh martensite (low IQ peak).
Fig. 4. Electron backscatter diffraction (EBSD) image quality (IQ) distributions of bcc phase in the Base and Hi-Al steels annealed at 700 and 725 °C for 1 h followed by water quenching. The intensities of the peaks at lower IQ and higher IQ reflect the fractions of martensite and ferrite, respectively.
Fig. 5. Equilibrium (a) C, (b) Mn and (c) Al contents in the intercritical austenite for the Base, Lo-Al and Hi-Al steels as a function of annealing temperature predicted by the Thermo-Calc® TCFE 7 database. The critical temperatures (A1, M0, ACM and A3) are shown (only) for the Hi-Al alloy.
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Due to its tetragonality and high dislocation density, the martensite exhibits relatively low IQ values compared to ferrite. For each alloy, the relative intensity of the lower IQ peak decreases at lower annealing temperature. In particular, the lower IQ peak almost disappears in the Hi-Al 700 °C sample, implying less martensite formation during quenching. In both steels, martensite fractions (i.e., magnitude proportional to peak height) increase with annealing temperature, consistent with the predicted results shown in Fig. 2. The retained austenite stabilities for the three alloys with different annealing temperatures were estimated from the retained austenite chemical compositions. Calculated equilibrium amounts of C, Mn and Al in the intercritically annealed austenite for the three alloys are plotted against the intercritical annealing temperature in Fig. 5a to c, assuming full partitioning. The four critical temperatures provided in each figure are for the Hi-Al alloy. With decreasing annealing temperature, Fig. 5a shows that the predicted austenite carbon content increased below A3, showed a maximum at ACM and decreased due to cementite precipitation below ACM. Equilibrium partitioning of Mn (Fig. 5b) is clearly greater at lower annealing temperatures where less austenite and more ferrite are present. The austenite Al content correspondingly decreases at lower intercritical annealing temperatures (Fig. 5c). To estimate the mechanical stability of the retained austenite for the three alloys, the MS temperatures and kS values were calculated using Eqs. (4) and (6) based on the chemical compositions predicted at each annealing temperature, and the results are shown in Fig. 6. Over the entire intercritical annealing temperature range, the MS temperatures of the three alloys gradually decrease with decreasing annealing temperature (Fig. 6a). Increased austenite stabilization originates from the
Fig. 6. Predicted (a) MS temperatures and (b) kS values of the retained austenite in the Base, Lo-Al and Hi-Al steels after intercritical annealing.
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enrichment of C and Mn and depletion of Al. The austenite stability exhibits increased temperature dependence in the Hi-Al alloy. Based on the MS temperature in Fig. 6a, Fig. 6b shows the calculated kS values using Eq. (4). Similar to the MS temperature, the kS values decrease with lower annealing temperature and Al additions amplified the reduction. Lower kS values imply less martensite transformation with strain. Fig. 7a to c show the measured room temperature tensile flow curves for the Base, Lo-Al and Hi-Al steels annealed at temperatures ranging from 550 to 725 °C for 1 h followed by water quenching. In each alloy, tensile deformation characteristics are strongly influenced by the annealing temperature. For example, the total elongation in the Base alloy annealed at 550 °C is 4.5%, and greatly increased up to 38.8% for the sample annealed at 650 °C. In addition, the maximum UTS value of 1603 MPa obtained at 725 °C is twice as large as the
Fig. 7. Room temperature engineering stress-strain curves of the (a) Base, (b) Lo-Al and (c) Hi-Al steels annealed at temperatures ranging from 550 to 725 °C for 1 h followed by water quenching.
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minimum value of 730 MPa obtained at 625 °C. As expected, samples annealed at temperatures between 650 and 700 °C having higher fγR values exhibited relatively larger elongations, and greater UTS was observed in the samples annealed above 700 °C having higher martensite fractions. The slopes of the tensile flow curves tend to increase with annealing temperature, implying increased strain hardening, reflecting increased transformation of austenite to martensite with strain. This behavior is consistent with predictions in Fig. 6b, where kS values increased with temperature, indicating more rapid martensite transformation during strain for microstructures processed at higher annealing temperatures, providing increased strain hardening. Similar evolutions of tensile flow curves with annealing temperatures are observed in both the Lo-Al (Fig. 7b) and the Hi-Al alloys (Fig. 7c). Experimental and predicted UTS and UE values as a function of annealing temperature for the Base, Lo-Al and Hi-Al steels are provided in Fig. 8. Fig. 8a shows that with an increase in annealing temperature, experimentally measured UTS values for the three alloys initially decrease, exhibit a minimum at temperatures near 650 °C, and dramatically increase at higher temperatures. For all tested conditions, Al additions decrease the predicted UTS values and the effects of Al additions are greater at higher annealing temperatures. Fig. 8b shows the corresponding UTS values predicted for the Base, Lo-Al and Hi-Al steels. The predictions exhibit similar variations with annealing temperatures as the experimental data. Predicted temperatures corresponding to the lowest UTS values in the Base, Lo-Al and Hi-Al steels are 610, 632 and 666 °C, respectively, which are equal or slightly less than the M0 temperature values (614, 632 and 666 °C for the Base, Lo-Al and Hi-Al steels, respectively), below which martensite formation is not anticipated during quenching from the intercritical heat treating temperature.
In Fig. 8c, experimentally measured UE values for the Base, Lo-Al and Hi-Al steels exhibit a peak as a function of annealing temperature. Al additions decrease the maximum UE values from 35.2 to 25.6% and shift the peaks to higher temperatures. Predicted UE values for the Base, Lo-Al and Hi-Al steels shown in Fig. 8d exhibit the maximum UE at 620, 640 and 670 °C, respectively. These temperatures are slightly higher than the temperatures for the lowest UTS. Overall, the predicted UE values with annealing temperature and Al additions show reasonable agreement with the experimental results. However, Al additions increase the maximum predicted UE value from 20.3, 22.1 to 25.3%, in contrast to the experimental results. One of the possible reasons for this discrepancy is the composition dependent phase tensile properties. While not considered for the predictions in the present study, tensile properties of each phase are influenced by their chemical compositions. Fig. 9a shows the relationship between UTS×UE values and annealing temperature for the Base, Lo-Al and Hi-Al steels obtained from predictions and experiments. Such products are often used to compare the overall strength-ductility balance for different microstructures. Overall, experimental UTS×UE values show higher maximum values and larger variations with annealing temperature compared to the predicted results. The experimental UTS×UE values for the Base, Lo-Al and Hi-Al alloys exhibit a maximum at 675, 675 and 700 °C, respectively, and predicted maximum values are obtained at 645, 654 and 671 °C (corresponding to the ACM for each alloy), respectively. For both experimental and predicted results, the maximum UTS×UE values are obtained at temperatures having the maximum fγR values. The relationships between the UTS×UE and fγR values for the three alloys are shown in Fig. 9b. While both predicted and experimental results show a strong relationship between the UTS×UE and fγR values, the effect of the fγR
Fig. 8. Tensile properties of the Base, Lo-Al and Hi-Al steels: (a) experimentally measured ultimate tensile strength (UTS, Exp.), (b) predicted ultimate tensile strength (UTS, Pre.), (c) experimentally measured uniform elongation (UE, Exp.), and (d) predicted uniform elongation (UE, Pre.).
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on the UTS×UE is somewhat underestimated in the predicted results. Also, the results in Fig. 9b illustrate the importance of retained austenite on the tensile properties of intercritically annealed medium Mn steels. Medium Mn steels after intercritical annealing have diverse combinations of various phases and austenite stability, resulting in large variations of tensile deformation behavior. The model introduced in the present study provides an important means of predicting tensile properties of medium Mn steels using predicted phase fractions and TRIP behavior using only thermodynamic inputs, without requiring any direct experimental observations. Through this modeling method, it is possible to estimate the complex effects of alloy composition and annealing temperature on the tensile deformation behavior of medium Mn steels. While the accuracy of the predicted results can be improved, the models should be useful in the design of alloy compositions and heat treatment processes for improved tensile properties of medium Mn steels. 5. Conclusions The tensile properties of intercritically annealed Base, Lo-Al and Hi-Al steels were predicted and the ultimate tensile strengths (UTS) and uniform elongations (UE) were compared to experimental data for three 0.19C, 4.5Mn steels with different Al additions. Intercritically annealed phase fractions and retained austenite chemical compositions were predicted using a retained austenite prediction model assuming ortho-equilibrium solute partitioning. The martensite transformation behavior during tensile deformation was estimated from the predicted retained austenite chemical composition. Tensile properties, i.e. UTS and UE of intercritically annealed medium Mn steels were predicted
Fig. 9. Ultimate tensile strength (UTS)-uniform elongation (UE) products for the Base, LoAl and Hi-Al steels obtained from experiments (Exp.) and predictions (Pre.) as a function of (a) intercritical annealing temperature and (b) retained austenite fraction (fγR).
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by incorporating the predicted phase fractions and martensite transformation kinetics in a simplified composite model for multi-component systems. Predictions show minimum UTS values near the M0 temperature, maximum UE values at temperatures between the M0 and ACM, and maximum UTS×UE values at ACM where the greatest retained austenite fractions were predicted. Overall, the evolution of predicted tensile properties with annealing temperature and Al additions exhibited reasonable agreement with the experimental data, reflecting the importance of retained austenite, and austenite stability. Acknowledgement The authors gratefully acknowledge the support of the sponsors of the Advanced Steel Processing and Products Research Center, an industry/university cooperative research center at the Colorado School of Mines. References [1] D.K. Matlock, J.G. Speer, E. De Moor, P.J. Gibbs, Recent developments in advanced high strength sheet steels for automotive applications: an overview, Jestech 15 (2012) 1–12. [2] I. Tamura, Deformation-induced martensitic transformation and transformationinduced plasticity in steels, Met. Sci. 16 (1982) 245–253. [3] R.L. Miller, Ultrafine-grained microstructures and mechanical properties of alloy steels, Metall. Trans. A 3 (1972) 905–912. [4] Y.-K. Lee, J. Han, Current opinion in medium manganese steel, Mater. Sci. Technol. 31 (2015) 843–856. [5] S. Lee, B.C. De Cooman, On the selection of the optimal intercritical annealing temperature for medium Mn TRIP steel, Metall. Mater. Trans. A 44 (2013) 5018–5024. [6] D.-W. Suh, S.-J. Park, T.-H. Lee, C.-S. Oh, S.-J. Kim, Influence of Al on the microstructural evolution and mechanical behavior of low-carbon, manganese transformationinduced-plasticity steel, Metall. Mater. Trans. A 41 (2010) 397–408. [7] T. Furukawa, Dependence of strength-ductility characteristics on thermal history in low carbon, 5 wt-%Mn steels, Mater. Sci. Technol. 5 (1989) 465–470. [8] H. Huang, O. Matsumura, T. Furukawa, Retained austenite in low carbon, manganese steel after intercritical heat treatment, Mater. Sci. Technol. 10 (1994) 621–626. [9] D.E. Kim, Y.K. Park, O.Y. Lee, K.G. Jin, S.J. Kim, Formation of retained austenite and mechanical properties of 4–8%Mn hot rolled TRIP steels, Korean J. Mater. Res. 15 (2005) 115–120. [10] P.J. Gibbs, E. De Moor, M.J. Merwin, B. Clausen, J.G. Speer, D.K. Matlock, Austenite stability effects on tensile behavior of manganese-enriched-austenite transformationinduced plasticity steel, Metall. Mater. Trans. A 42 (2011) 3691–3702. [11] D.W. Suh, J.H. Ryu, M.S. Joo, H.S. Yang, K. Lee, H.K.D.H. Bhadeshia, Medium-alloy manganese-rich transformation-induced plasticity steels, Metall. Mater. Trans. A 44 (2013) 286–293. [12] D.K. Matlock, J.G. Speer, Design considerations for the next generation of advanced high strength steels, in: H.C. Lee (Ed.), Proceedings of the 3rd International Conference on Advanced Structural Steels, Kor. Inst. Metal. Mater., Seoul, South Korea 2006, pp. 774–781. [13] R. Rana, P.J. Gibbs, E. De Moor, J.G. Speer, D.K. Matlock, A composite modeling analysis of the deformation behavior of medium manganese steels, Steel Res. Int. 85 (2015) 1–12. [14] S. Lee, Y. Estrin, B.C. De Cooman, Constitutive modeling of the mechanical properties of V-added medium manganese TRIP steel, Metall. Mater. Trans. A 44 (2013) 3136–3146. [15] Y. Estrin, H. Mecking, A unified phenomenological description of work hardening and creep based on one-parameter models, Acta Mater. 32 (1984) 57–70. [16] E. De Moor, D.K. Matlock, J.G. Speer, M.J. Merwin, Austenite stabilization through manganese enrichment, Scr. Mater. 64 (2011) 185–188. [17] S. Kang, E. De Moor, J.G. Speer, Retained austenite stabilization through solute partitioning during intercritical annealing in C-, Mn-, Al-, Si-, and Cr-alloyed steels, Metall. Mater. Trans. A 46 (2015) 1005–1011. [18] S. Kang, E. De Moor, J.G. Speer, D. Krizan, in: M. Militzer, G. Botton, L.-Q. Chen, J. Howe, C. Sinclair, H. Zurob (Eds.),Aluminum alloyed medium manganese steels, Proceedings of the International Conference on Solid-Solid Phase Transformations in Inorganic Materials 2015 2015, pp. 1145–1151 (Whistler, BC, Canada). [19] ASTM E8/E8M-11, Standard Test Methods for Tension Testing of Metallic Materials, West Conshohocken, PA, ASTM International, 2011. [20] C.F. Jatczak, Retained austenite and its measurement by X-ray diffraction, SAE Technical Paper No. 800426, 1980. [21] S.T. Mileiko, The tensile strength and ductility of continuous fibre composites, J. Mater. Sci. 4 (1969) 974–977. [22] G. Garmong, R.B. Thompson, A theory for the mechanical properties of metal-matrix composites at ultimate loading, Metall. Trans. 4 (1973) 863–873. [23] R.G. Davies, The deformation behavior of a vanadium-strengthened dual-phase steel, Metall. Trans. A 9A (1978) 41–52. [24] G.R. Speich, R.L. Miller, Mechanical properties of ferrite–martensite steels, in: R.A. Kot, J.W. Morris (Eds.), Structure and Properties of Dual-phase Steels, TMS, Warrendale, PA 1979, pp. 145–182.
146
S. Kang et al. / Materials and Design 97 (2016) 138–146
[25] J. Becker, E. Hornbogen, F. Wendl, Analysis of mechanical properties of a low-alloyed mo-steel with different phosphorus additions, in: R.A. Kot, B.L. Bramfitt (Eds.), Fundamentals of Dual-Phase Steels, TMS, Warrendale PA 1981, pp. 383–398. [26] S. Lee, Y. Estrin, B.C. De Cooman, Constitutive modeling of the mechanical properties of V-added medium manganese TRIP steel, Metall. Mater. Trans. A 44A (2013) 3136–3146. [27] H.N. Han, C.S. Oh, G. Kim, O. Kwon, Design method for TRIP-aided multiphase steel based on microstructure-based modelling for transformation induced plasticity and mechanically induced martensitic transformation, Mater. Sci. Eng. A 499 (2009) 462–468. [28] O. Bouaziz, J.D. Embury, Microstructural design for advanced structural steels, Mater. Sci. Forum 539-543 (2007) 42–50. [29] K.S. Choi, A. Soulami, W.N. Liu, X. Sun, M.A. Khaleel, Influence of various material design parameters on deformation behaviors of TRIP steels, Comput. Mater. Sci. 50 (2010) 720–730.
[30] G.E. Dieter, Mechanical Metallurgy 3rd Edition, McGraw-Hill, London, 1986. [31] N. Tsuchida, H. Masuda, Y. Harada, K. Fukaura, Y. Tomota, K. Nagai, Effect of ferrite grain size on tensile deformation behavior of a ferrite-cementite low carbon steel, Mater. Sci. Eng. A 488 (2008) 446–452. [32] A. Clarke, Carbon Partitioning Into Austenite From Martensite in a Silicon-containing High Strength Sheet Steel(PhD Thesis) Colorado School of Mines, 2006. [33] W.A. Poling, Grain Size Effects in Micro-tensile Testing of Austenitic Stainless Steel(MS Thesis) Colorado School of Mines, 2012. [34] K.W. Andrews, Empirical formulae for the calculation of some transformation temperatures, J. Iron Steel Inst. 203 (1965) 721–727. [35] K. Sugimoto, M. Kobayashi, S. Hashimoto, Ductility and strain-induced transformation in a high-strength transformation-induced plasticity-aided dual-phase steel, Metall. Mater. Trans. A 23 (1992) 3085–3091.