# Effects of Partial Replacement of Si by Al on Cold Formability in Two Groups of Low-Carbon Third-Generation Advanced High-Strength Steel Sheet: A Review

## Abstract

**:**

## 1. Introduction

_{R}), reverted austenite, and/or austenite. The representative characteristic is shown by the product of tensile strength and total elongation (TS×TEl), which increases with an increase in the initial volume fraction of austenite or retained austenite (fγ

_{0}) (Figure 1) [6]. For a TS level higher than 1.0 GPa, good performance and low production costs can be obtained by using the third-generation AHSSs.

**Figure 1.**Relationship between the product of tensile strength and total elongation (TS×TEl) and an initial volume fraction for retained austenite, reverted austenite, or austenite (fγ

_{0}) in the first-, second-, and third-generation (Group I and Group II) advanced high-strength steel (AHSS) sheets. Q&T: quenching and tempering martensitic steel, DP: dual-phase steel, CP: complex-phase steel, TPF, TAM, TBF, and TM: transformation-induced plasticity (TRIP)-aided steels with polygonal ferrite, annealed martensite, bainitic ferrite, and martensite matrix structure, respectively. Q&P: one-step and two-step quenching and partitioning steels, CFB: carbide-free bainitic steel, D-MMn: duplex type medium-Mn steel, L-MMn: laminate type medium-Mn steel, BF-MMn: bainitic ferrite-type medium-Mn steel, Q&P-MMn: Q&P-type medium-Mn steel, M-MMn: martensite-type medium-Mn steel, HMn TWIP: high-manganese TWIP steel, Aus: austenitic steel. This figure is reproduced based on references [6,37].

_{R}during the heat treatment, in the same way as P [66,67]. As Al does not deteriorate the coatability (or galvanizing property), unlike Si [48,49,68,69,70], it becomes especially advantageous for industrial production in conventional galvanizing lines. However, Al is a weak solid-solution-strengthening element in steel. In addition, the content is limited because Al is a ferrite-stabilizing element [69].

## 2. Microstructure and Retained Austenite Characteristics

#### 2.1. C-Si-Mn steel

_{IT}) above M

_{s}(IT process (i)) or between M

_{s}and M

_{f}(IT process (ii)) are conducted. The CFB [27,28,29,30,31,32,33,34] and BF-MMn [45] steels are produced by the same heat treatment as the TBF steel (Figure 2a). The one-step Q&P steel [20] involves the IT process (ii) of the TBF steel (Figure 2b). For the TM steel, the IT process below M

_{f}(iii) is applied after austenitizing (Figure 2a) [52]. The heat treatment corresponding to the IT process (iii) contains direct quenching to room temperature and subsequent partitioning (DQ&P) process. The M-MMn steel is fabricated by the same IT process (iii) as the TM steel [57,58,59].

_{bf}) for the IT process (i) (Figure 3a) [81], and it a mixture of α

_{bf}and primary coarse soft martensite (α

_{m}) for the IT process (ii) (Figure 3b) [81]. In the TBF steel, an initial γ

_{R}fraction (fγ

_{0}) increases with increasing T

_{IT}(Figure 4) [51]. The highest initial carbon concentration (Cγ

_{0}) is obtained in the TBF steel subjected to the IT process at the temperatures between M

_{s}and M

_{f}. In the TBF steel subjected to the IT process (ii), a small amount of MA phase (a mixture of secondary fine hard martensite (α

_{m}

^{*}) and film-like γ

_{R}) exists as the second phase. In addition, only a small amount of carbide (θ) precipitates only in the α

_{m}lath structure [15]. The microstructure of CFB [27,28,29,30,31,32,33,34], one-step Q&P [20], and BF-MMn [45] steels resembles that of TBF steels subjected to the IT processes (i) and/or (ii). In general, the above-mentioned fγ

_{0}and Cγ

_{0}are calculated by the methods proposed by Maruyama [82] and Dyson and Holmes [83], respectively.

_{m}or auto-tempered primary martensite (Figure 3c) [81]. The TM steel contains a large amount of MA phase and a small amount of θ in the α

_{m}lath structure as the second phase [51,52,53]. The θ fraction (f

_{θ}) increases with decreasing T

_{IT}[52], although it is much lower than that of quenching and tempering (Q&T) steel [52,54]. The TM steel is also called “TRIP-aided duplex martensitic steel” because its microstructure consists of α

_{m}matrix structure and a large amount of MA second phase. It is noteworthy that M-MMn steel contains much larger amounts of MA phase and γ

_{R}than TM steel [59].

_{Q}) between M

_{S}and M

_{f}after austenitizing and subsequent partitioning at a temperature (T

_{P}) higher than M

_{s}(Figure 2b) [19,20,48]. The process forms the microstructure of α

_{bf}and α

_{m}matrix and γ

_{R}, similar to the IT process (ii) for TBF and CFB steels and the one-step Q&P process (Figure 3b). The variation in volume fractions of various phases as a function of T

_{Q}is illustrated in Figure 5a. On the quenching to T

_{Q}, a certain amount of austenite transforms to α

_{m}first. The α

_{m}fraction (fα

_{m}) can be estimated by the following empirical equation proposed by Koistinen and Marburger [84].

_{m}= 1 − exp {−1.1 × 10

^{−2}(M

_{S}− T

_{Q})}

_{Q}is close to M

_{f}, a small amount of carbide (θ) precipitates only in the α

_{m}lath structure [23,49]. During subsequent partitioning at temperatures above M

_{s}, most of the remaining austenite transforms into α

_{bf}. At the same time, the α

_{m}softens through carbon migration (carbon enrichment) into untransformed austenite and carbide precipitation [23]. During final cooling to room temperature, a part of unstable austenite transforms into the MA phase. Typical two examples of T

_{Q}dependence of fγ

_{0}in two-step Q&P steels with different carbon content are shown in Figure 5b. In these two-step Q&P steels, the optimum T

_{Q}which gives the maximum volume fraction of γ

_{R}is between M

_{s}and M

_{f}[19,20,23].

_{R}. The mechanical stability is mainly related to its Cγ

_{0}, along with the size and stacking fault energy (SFE) of γ

_{R}, matrix structure, deformation temperature, other microalloying elements, etc. [85]. According to Sugimoto et al. [86], the mechanical stability of γ

_{R}can be defined by the following k-value (or the strain-induced transformation factor),

_{0}− ln fγ)/ε

_{0}− ln fγ = k

_{1}ΔG

^{α’γ}ε

_{1}is a modified k-value. ΔG

^{α’γ}(= G

^{α’}− G

^{γ}) is the chemical free-energy change for the transformation of austenite to ferrite (martensite) with the same composition (without considering stored energy due to the shape deformation), where G

^{α’}and G

^{γ}are the chemical free energies of ferrite (martensite) and austenite, respectively.

_{1}-values decrease with increasing Cγ

_{0}in the TPF steel. As shown in Figure 6 [88], the k-values of low-carbon Si-Mn TBF, TM, D-MMn, and M-MMn steels also decrease with increasing Cγ

_{0}, although the k-values are about three times higher than those of TPF and TAM steels because of lower Cγ

_{0}[89]. It is noteworthy that the k-values of 3Mn and 5Mn M-MMn steels are lower than those of TBF and TM steels, like 3Mn D-MMn steel. This is because high solute Mn concentration in the γ

_{R}plays a role in the austenite stabilizer.

#### 2.2. C-Si/Al-Mn Steel

_{0}) at which austenite and martensite have the same chemical free energy in steel [85]. Only Co increases the T

_{0}and makes the austenite unstable. Many ferrite-stabilizing elements also stabilize the austenite, although they increase the T

_{0}. Exceptionally, Cr lowers the T

_{0}and significantly increases the austenite stability by the addition of several mass% points, although it is a ferrite-stabilizing element [85].

_{0}curves calculated for 19C-1.54Si-1.51Mn-0.04Al (0Al), 0.5Al: 0.20C-0.99Si-1.51Mn-0.49Al (0.5Al), and 0.20C-0.49Si-1.50Mn-0.99Al (1.0Al) TPF steels are shown in Figure 7b [75]. Al shifts the T

_{0}line to the high-carbon-concentration side. This means that Al also plays a role in lowering the k-value or increasing the mechanical stability [88]. Al increases the SFE of γ

_{R}, and Cu and Si [92], which makes TRIP and TWIP difficult [93]. However, there are some disadvantages to adding Al: it can reduce solid-solution-strengthening and raises the M

_{s}[91].

_{R}and decreases its volume fraction. In this case, the increased mechanical stability is mainly associated with higher Cγ

_{0}and higher SFE. A similar effect of Al is obtained in 0.2C-(0.2-1.5)Si-1.24Mn-(0.02-1.22)Al-0.2Cr-0.003B TM steels [55,88] (Figure 6), although the mechanical stability and volume fraction are lower than those of 0.2C-(0.5-1.5)Si-1.5Mn-(0.04-1.0)Al-(0-0.2)Mo-(0-0.05)Nb [12] and 0.2C-(1.0-1.544)Si-1.5Mn-(0.04-0.5)Al-(0-0.05)Nb TBF steels [14] and 0.2C-(0.5-1.5)Si-1.5Mn-(0.04-1.0)Al TPF and TAM steels [75,78]. According to Sugimoto et al. [55], this is caused by insufficient carbon enrichment during the IT process at a lower temperature after the DQ process, which leads to a large amount of MA phase. Additionally, a similar effect of Al on the k-value or Cγ

_{0}has been reported for 0.2C-1.8Si-2Mn-(0-0.5)Al-0.2Mo-1.0Cr CFB [32], 0.2C-8Mn-(0-3)Al D-MMn [42], and 0.2C-(0.08-1.5)Si-4Mn-(0.02-1.46)Al Q&P-MMn [49] steels.

_{bf}lath, and film-like

**γ**in 0.2C-(0.5-1.5)Si-1.5Mn-(0.04-1.0)Al-(0-0.2)Mo-(0-0.05)Nb [12] and 0.2C-(0.5-1.5)Si-(1.5-2.5)Mn-(0.04-1.0)Al TBF steels [13]. Zhu et al. [28] show that Al addition of 1.0 to 1.5 mass% results in a remarkable refinement of α

_{R}_{bf}lath, film-like

**γ**, and MA island in 0.25C-(0.1-1.09)Si-2.07Mn-(0.02-1.54)Al CFB steels. A similar result was also found by Tian et al. [32]. For θ precipitation, He et al. [31] found that Al addition successfully suppresses the formation of θ in 0.25C-2.07Mn-(0.02-1.54)Al CFB steels through effective carbon enrichment from transformed α

_{R}_{bf}to the adjacent untransformed austenite [94]. Kaar et al. [47] and Wallner et al. [49] report that a significantly larger amount of triaxial aligned θ is precipitated in the α

_{m}matrix in 0.173C-4.46Mn-1.47Si-0.03Al and 0.20C-4.52Mn-0.04Si-1.31Al two-step Q&P-MMn steels, respectively, although Al exhibits lower suppression of θ precipitation than Si.

_{m}transformation, Kobayashi et al. [55] found that the α

_{m}size is largely unchanged by Al addition in 0.2C-(0.2-1.5)Si-1.24Mn-(0.02-1.22)Al-0.2Cr-0.003B TM steels, unlike the above results relating to the α

_{bf}transformation. In this case, the prior austenitic grain size was nearly the same in both sheets of steel. Kantanen et al. [24] also showed that the α

_{m}size is not influenced by Al content in 0.30C-0.56Si-2.00Mn-1.10Al-2.20Cr two-step Q&P steel.

## 3. Tensile Properties

#### 3.1. C-Si-Mn Steel

_{R}of 4 to 30 vol% is formulated by

^{M}(ε) + Δσ

_{h}(ε)

^{M}(ε) and Δσ

_{h}(ε) are the flow stress of the matrix and strain hardening increment of the steel, respectively. The Δσ

_{h}(ε) can be estimated by

_{h}(ε) = Δσ

_{i}(ε) + Δσ

_{t}(ε) + Δσ

_{f}(ε)

_{i}(ε), Δσ

_{t}(ε), and Δσ

_{f}(ε) represent “the long-range internal stress hardening”, “the strain-induced transformation hardening”, and “the forest dislocation hardening”, respectively, which can be formulated by

_{i}(ε) = {(7−5ν)μ/5(1-ν)} f·ε

_{p}

^{u}

_{t}(ε) = g(Δfα

_{m})

_{f}(ε) = ζμ (

**b**·f·ε/2r)

^{1/2}

_{p}

^{u}is “the eigenstrain” [95], f is the volume fraction of the second phase, g(Δfα

_{m}) is a function of the strain-induced martensite fraction, ζ is a material constant,

**b**is the Burgers vector, and r is particle radius of the second phase. In the D-MMn steel, the second phase is mainly film-like γ

_{R}. In the TBF, CFB, one-step Q&P, and BF-MMn steels subjected to the IT process at the temperatures above M

_{s}, the second phase corresponds to untransformed carbon-enriched γ

_{R}and strain-induced martensite. In these steels subjected to the IT process at the temperatures between M

_{s}and M

_{f}, the second phases correspond to α

_{m}, untransformed γ

_{R}, and strain-induced martensite, in the same way as two-step Q&P and two-step Q&P-MMn steels. Sometimes, a small amount of MA phase is also classified into the second phase in these steels. On the other hand, the second phase of the TM and M-MMn steels is mainly equivalent to the MA phase. The strain-hardening mechanism of two-step Q&P steel is also suggested by Celada-Casero et al. [25].

_{Q}, T

_{IT}, T

_{P}, and these holding times. The effects of T

_{Q}on the tensile properties in (0.2–0.3)C-1.6Si-4.0Mn-1.0Cr steels subjected to the two-step Q&P process are shown in Figure 8 [23]. The largest TEls of both sheets of steel were obtained by quenching at T

_{Q}= M

_{s}− 100 °C, which was about 50 °C lower than T

_{Q}for the largest volume fraction of γ

_{R}(Figure 4b). The optimum T

_{Q}s for TEl roughly match those for the minimum TS. On the other hand, the minimum yield stress (YS) was about 50 °C higher than the optimum T

_{Q}for the TEl.

_{IT}on the tensile properties of 0.20C-1.50Si-1.50Mn-0.05Nb TBF and TM steels are shown in Figure 9 [51]. In these steels, the largest UEL, TEL, and TS×TEl can be obtained at T

_{IT}above M

_{s}, although the YS and TS considerably decrease. The optimum T

_{IT}for tensile ductility agrees well with one for the largest fraction of γ

_{R}[51]. In the TM steel subjected to the IT process at temperatures below M

_{f}, relatively high TS and TS×TEl are achieved, compared to the TBF steels subjected to the IT process at temperatures between M

_{s}and M

_{f}. In this case, the YS of the TM steel decreases due to a large amount of MA phase, resulting in the continuous yielding [96], in the same way as ferrite-martensite DP steel.

_{R}or small Δfα

_{m}.

#### 3.2. C-Si/Al-MnSsteel

_{s}and M

_{f}[14]. If the addition of Al is below 0.7 mass% in the TBF and TM steels, these steels keep the TS×TEl of 10 GPa%, which is slightly lower than those of Al-free TBF and TM steels [14,88] (Figure 10b). However, Al addition of 1.2 mass% (or the partial replacement of Si by 1.2 mass% Al) considerably decreases the TS×TEl in TM steel in a TS range above 1.2 GPa (see 1.2Al TM steel in Figure 10b).

_{m}lath structure hardly influences the ductility of the Al-added TM steels because the θ fraction is very little. According to Jing et al. [42], 1.39 mass% Al addition achieved the highest TS×TEl in (0.18-0.19)C-(7.66-7.93)Mn-(0-2.79)Al D-MMn steels.

_{R}[14,39,40,51,53,55,59,88]. In addition, they increase as the k-value decreases (Figure 11b). Therefore, a decrease in TS×TEl of 1.2Al TM steel may be caused by the decreased volume fraction of γ

_{R}and low solid solution-hardening, although the mechanical stability of γ

_{R}increases (or the k-value decreases). The TS×TEl− fγ

_{0}relation of 5Mn M-MMn steel is superior to those of the other TBF and TM steels. This is associated with the increased mechanical stability (decreased k-value), the volume fraction of γ

_{R}, and the MA phase due to its high Mn concentration [88]. 5Mn D-MMn steel has a higher TS×TEl than 5Mn M-MMn steel. This is mainly caused by the high-volume fraction of γ

_{R}.

## 4. Stretch formability

#### 4.1. C-Si-Mn Steel

_{max}), which is usually measured using a ball-head punch. Figure 12a shows the T

_{IT}dependence of the H

_{max}in 0.20C-1.50Si-1.50Mn-0.05Nb TBF and TM steels subjected to the IT process at T

_{IT}= 200 °C to 450 °C for 200 s after austenitizing [51]. For comparison, the T

_{P}dependence of the H

_{max}of 0.20C-1.50Si-1.50Mn-0.05Nb TM steel subjected to the DQ&P process after austenitizing and then partitioning at T

_{P}= 200 °C to 450 °C for 1000 s is shown in the Figure 12a. The H

_{max}s values of the IT-processed TBF and TM steels (red line) increase with increasing T

_{IT}. The H

_{max}of DQ&P-processed TM steel increases with increasing T

_{P}, although the H

_{max}s are lower than those of IT-processed TBF and TM steels. This result indicates that the IT process is suitable compared to the DQ&P process in TBF and TM steels. According to Kobayashi et al. [51], this result is associated with a larger amount of γ

_{R}. As shown in Figure 12b, the products of TS and H

_{max}(TS×H

_{max}s) of the IT-processed TBF and TM steels and DQ&P-processed TM steel are much higher than those of 0.2C-1.5Si-(1.5-5.0)Mn M-MMn steels [59], 0.082C-0.88Si-2.0Mn ferrite-martensite DP steel [16,51,52], and 0.23C-0.19Si-1.29Mn-0.21Cr-0.003B 22MnB5 steel subjected to the Q&T process (22MnB5 Q&T steel) [16,52] in a TS range above 1.0 GPa, although they are a little lower than those of 0.2C-1.5Si-(1.5-5.0)Mn D-MMn steels [39]. It is very important to know that large tensile ductility does not necessarily lead to high stretch formability because of the different stress states, although it exhibits a linear relationship with stretch formability. Namely, equi-biaxial tension growing on stretch-forming promotes crack and/or void initiation compared to uniaxial tension [80].

_{max}of 5Mn D-MMn steel is further increased to 10.5 mm by warm forming at 200 °C, which significantly increases the mechanical stability of γ

_{R}[39].

#### 4.2. C-Si/Al-Mn Steel

_{max}decreases with increasing Al content in the same way as the TEl and TS×TEl in (Figure 13a) [88]. It is interesting that partial replacement of Si by 1.2 mass% Al considerably decreases the H

_{max}and TS×H

_{max}[55,88], compared to those of 0.2C-1.5Si-1.5Mn-0.05Nb (1.5Si) TBF steel [51] and 0.2%-1.5Si-1.5Mn-(0-1.0)Cr-(0-0.2)Mo (Cr-Mo) TM steels [53].

_{max}exhibits a linear relationship with the TEl in 0Al, 0.7Al, and 1.2Al TM steels, the Cr-Mo TM steels, and the 1.5Si TBF steels, although the slopes of 5Mn M-MMn and D-MMn steels are lower than those of these steels. Kobayashi et al. [55] propose that lower H

_{max}and TS×H

_{max}of 1.2Al TM steel may be caused by lower UEl and Tel, resulting in low solution-hardening and low γ

_{R}fraction. According to Sugimoto et al. [59], the small H

_{max}of the 5Mn M-MMn steel is caused by the presence of a much larger MA phase, although a large amount of metastable γ

_{R}makes a positive contribution to the H

_{max}. In this case, a crack initiation at the interface between the matrix and MA phase is promoted by equi-biaxial tension, as opposed to uniaxial tension (tensile test).

_{max}has not been investigated for the third-generation AHSSs. It is expected that the complex addition of Al and Nb enhances the H

_{max}because it achieves a large TEl in 0.5Al-0.05Nb TBF steel (Figure 10b).

## 5. Stretch-Flangeability

#### 5.1. C-Si-Mn Steel

_{f}− d

_{0})/d

_{0}

_{0}and d

_{f}are the original diameter of the punched hole and the hole diameter upon cracking during the hole-expansion test, respectively. In many cases, the hole-punching tests are carried out at a clearance of about 10% [98]. The subsequent hole-expansion tests are conducted using a conical punch tool with a vertical angle of 60 deg. [55,59]. Recently, the conical punch tool has been preferentially used to measure the HER.

_{IT}in 0.20C-1.50Si-1.50Mn-0.05Nb IT-processed TBF and TM steels and increases with increasing T

_{P}in DQ&P-processed TM steel (Figure 14a) [51]. In this case, the TM steel subjected to the IT process at the temperatures of M

_{f}(50 to 100) °C and TBF steel subjected to the IT process at the temperatures between M

_{s}and M

_{f}achieve higher TS×HER (50 to 60 GPa%) than the DQ&P-processed TM steel. In addition, the IT-processed TBF and TM steels have much higher TS×HER than 5Mn M-MMn [59] and 5Mn D-MMn steels [40], ferrite-martensite DP steel [16,51,52], and 22MnB5 Q&T steel [16,52] (Figure 14b). According to Kobayashi et al. [51], the high TS×HER of the IT-processed TM steel is mainly caused by (i) uniform α

_{m}lath structure with low θ fraction and (ii) plastic relaxation of localized stress concentration by the strain-induced transformation of metastable γ

_{R}at the α

_{m}lath boundary and/or in a finely dispersed MA phase, which suppresses the void and/or cracks initiation on punching and void coalescence or cracking on hole expansion. For the IT-processed TBF steel, the excellent stretch-flangeability is also associated with (iii) a uniform fine mixture of α

_{bf}and α

_{m}, and the above (i) and (ii). The effects of T

_{IT}on the HERs are also reported for 0.20C-1.40Si-1.70Mn-0.045Nb TBF steel [17], 0.2C-0.2Si-2Mn-0.03Ti-0.003B TBF steel [16], 0.13C-1.35Si-2.10Mn-0.98Cr (and 0.18C-1.40Si-2.13Mn-1.00Cr) TBF and DQ-processed TM steels [56], and 0.22C-1.48Si-3.79Mn-0.98Cr Q&P-MMn steel [46]. The effect of T

_{Q}(in the Q&P process) on the stretch-flangeability was reported by Im et al. [25] using 0.18C-1.5Si-2.6Mn steel (M

_{s}= 368 °C). In this case, higher HER was obtained by a lower quenching process at T

_{Q}(= 280 °C), followed by partitioning at T

_{P}= 425 °C.

#### 5.2. C-Si/Al-Mn Steel

_{m}lath structure, the high mechanical stability of γ

_{R}, and low solution-hardening.

_{s}, the partial replacement of Si by Al also achieved high TS×HER (see 0.5Al TBF steel in Figure 15b), in the same way as Al-free (0Al) TBF steel (see 0Al TBF steel in Figure 15b) [14]. Furthermore, the complex addition of 0.5 mass% Al and 0.05 mass% Nb considerably enhances the TS×HER (see 0.5Al-0.05Nb TBF steel in Figure 15b), compared to 0.5Al TBF steel. A similarly high TS×HER was also obtained by the complex addition of Al and Nb/Mo in 0.20C-(0.49-1.54)Si-(1.48-1.51)Mn-(0.04-0.99)Al-(0-0.05)Nb-(0-0.20)Mo TBF steels [12]. According to Sugimoto et al. [12,14], this increased stretch-flangeability by complex addition is mainly associated with the small punching damage due to refined microstructure, stabilized film-like γ

_{R}, and TRIP effect on hole expansion. In this case, the complex addition brings on precipitation-hardening by fine NbC/Mo

_{2}C.

## 6. Bendability

#### 6.1. C-Si-Mn Steel

_{min}) [52,99]. Figure 17a shows the variation in average bending angle as a function of austempering temperature (or T

_{IT}) in 0.34C-1.65Si-1.94Mn-1.07Cr CFB steel [29]. In this study, three-point bending tests were conducted to measure the average bending angle following the standard of the Association of German Automobile Industries (VDA 238-100). The largest average bending angle was obtained by the IT process at T

_{IT}= 350 °C, just higher than M

_{s}(329 °C), as shown in Figure 17a. The optimum T

_{IT}corresponds to that of maximum total elongation (TE) (Figure 17b). In this case, the microstructure was a mixture of α

_{bf}and γ

_{R}, and the γ

_{R}fraction was relatively high (Figure 17c). According to Rana et al. [29], the higher bendability is connected with a finer overall microstructure, the absence of α

_{m}, and the high mechanical stability of γ

_{R}.

_{min}as functions of the TS in 0.21C-1.49Si-1.50Mn-1.0Cr-0.05Nb TBF and TM steels, 0.082C-0.88Si-2.0Mn ferrite-martensite DP steel, and 22MnB5 Q&T steel [52]. Small R

_{min}was obtained in the TBF steel subjected to the IT process at the temperatures between M

_{s}and M

_{f}and in the TM steel subjected to the IT process at 200 °C (<M

_{f}= 261 °C). The high bendability of the TM steel is caused by uniform microstructure, resulting in high local ductility despite a high MA phase fraction [52].

_{min}s of the TPF, TAM, and TBF steels with the chemistry of (0.1-0.6)C-1.5Si-1.5Mn [99]. Small R

_{min}values are achieved in TBF and TAM steels with lath-type uniform microstructure and high mechanical stability of γ

_{R}, resulting in high local ductility. In parallel with the bendability, much research to improve the spring back is also conducted for applications of the third-generation AHSSs [100].

#### 6.2. C-Si/Al-Mn Steel

_{min}in the first-generation AHSSs was reported by Sugimoto et al. [75]. Figure 19 shows the effect of Al content on R

_{min}− T

_{IT}relation in 0.2C-(0.5-1.5)Si-1.5Mn-(0.04-1.0)Al TPF and TAM steels with TS between 700 and 1000 MPa [75]. Partial displacement by 0.5 and 1.0 mass% Al under a condition of Si+Al = 1.5 mass% reduced the R

_{min}in both sheets of steel. The optimum bendability was obtained in the TAM steel subjected to the IT process at temperatures between 275 and 425 °C, in the same way as the HER. According to Sugimoto et al. [75], the improved bendability of TAM steel is principally owed to the refined annealed martensite lath structure and the increased carbon concentration of γ

_{R}needles. In general, small R

_{min}is brought about by large RA. As Al-added TBF and TM steels have large RA values, a small R

_{min}is expected to be achieved in the Al-added TBF and TM steels.

## 7. Summary and Perspectives

_{R}and increases its mechanical stability in the third-generation AHSSs. In the TBF and TM steels, the partial replacement of Si by 0.7 mass% Al keeps the same large TEl and TS×TEl as Al-free steels. The TEL and TS×TEl decrease with increasing Al content, but the replacement of Si by 1.2 mass% Al deteriorates the TS×TEl in the TM steel with TS above 1.2 GPa, accompanied by decreases in YS, TS, and TEl. This is mainly caused by low solid solution-hardening and a decreased γ

_{R}fraction, resulting in decreased flow stress and a decreased strain hardening rate, despite increased mechanical stability of γ

_{R}. Similar results are reported for CFB, Q&P, and Q&P-MMn steels.

_{R}, despite low solution hardening and a decreased γ

_{R}fraction. Unfortunately, there is not any research on the bendability in Al-added AHSSs.

_{R}, and TRIP effect on hole-expanding. In this case, the complex addition brings on precipitation hardening by fine NbC/Mo

_{2}C. Further research on the cold formabilities of the AHSSs with a complex addition of Al and other elements (Nb, T, V, Cr, Mo, Ni, B, etc.), resulting in higher tensile strength above 1.5 GPa, is also expected in the future.

## Funding

## Institutional Review Board Statement

## Informed Consent Statement

## Conflicts of Interest

## Nomenclature

AHSS | advanced high-strength steel | TRIP | transformation-induced plasticity |

TWIP | twin-induced plasticity | HMn TWIP | high Mn TWIP |

Aus. | austenitic | TBF | TRIP-aided bainitic ferrite |

Q&P | quenching and partitioning | CFB | carbide-free bainite |

D-MMn | duplex type medium Mn | L-MMn | laminate type medium Mn |

BF-MMn | bainitic ferrite type medium Mn | Q&P-MMn | Q&P type medium Mn |

TM | TRIP-aided martensite | M-MMn | martensite type medium Mn |

TPF | TRIP-aided polygonal ferrite | TAM | TRIP-aided annealed martensite |

DP | dual-phase | CP | complex phase |

Q&T | quenching and tempering | DQ&P | direct quenching and partitioning |

IT | isothermal transformation | M_{s} | martensite-start temperature |

M_{f} | martensite-finish temperature | T_{IT} | isothermal transformation temperature |

T_{Q} | quenching temperature | T_{P} | partitioning temperature |

T_{0} | critical temperature at which austenite and martensite have the same chemical free energy | ||

γ_{R} | retained austenite | α_{bf} | bainitic ferrite |

α_{m} | primary coarse soft martensite | α_{m}* | secondary fine hard martensite |

MA | MA (α_{m}^{*}+γ_{R}) phase | θ | carbide |

fγ_{0} | initial volume fraction of γ_{R} | fγ | volume fraction of γ_{R} |

fα_{bf} | bainitic ferrite fraction | fα_{m} | primary martensite fraction |

fα_{m}* | secondary martensite fraction | f_{MA} | MA phase fraction |

fα_{m}’ | fα_{bf} + fα_{m}* | f_{θ} | carbide fraction |

Cγ_{0} | initial carbon concentration of γ_{R} | ε | plastic strain |

ΔG^{α’γ} | chemical free energy change for transformation of γ to α | G^{α’} | chemical free energy of ferrite (martensite) |

G^{γ} | chemical free energy of austenite | k | strain-induced transformation factor |

k_{1} | modified k-value | SFE | stacking fault energy |

σ | flow stress of steel | σ^{M} | flow stress of matrix |

Δσ_{h} | strain hardening increment | Δσ_{i} | long-range internal stress |

Δσ_{t} | transformation hardening | Δσ_{f} | forest dislocation hardening |

ν | Poisson’s ratio | μ | Shear modulus |

f | volume fraction of second phase | ε_{p}^{u} | eigenstrain |

Δfα_{m} | strain-induced martensite fraction | ζ | material constant |

b | Burgers vector | r | particle radius of second phase |

YS | yield stress | TS, UTS | tensile strength |

UEl | uniform elongation | TEl, TE | total elongation |

RA | reduction of area | TS×TEl | product of TS and TEl |

H_{max} | maximum stretch height | TS×H_{max} | product of TS and H_{max} |

HER | hole expansion ratio | TS×HER | product of TS and HER |

ss/t | shear section length to sheet thickness | R_{min} | minimum bending radius |

## References

- Rana, R.; Singh, S.B. Automotive Steels—Design, Metallurgy, Processing and Applications; Woodhead Publishing: Cambridge, UK, 2016; pp. 1–469. [Google Scholar]
- Fan, D.; Fonstein, N.; Jun, H. Effect of microstructure on tensile properties and cut-edge formability of DP, TRIP, Q&T and Q&P steels. AIST Trans.
**2016**, 13, 180–184. [Google Scholar] - Bleck, W.; Guo, X.; Ma, Y. The TRIP effect and its application in cold formable sheet steels. Steel Res. Int.
**2017**, 88, 1700218. [Google Scholar] [CrossRef] - Krizan, D.; Steinder, K.; Kaar, S.; Hebesberger, T. Development of third generation advanced high strength steels for automotive applications. In Proceedings of the 19th International Scientific Conference Transfer 2018, Trencin, Slovakia, 22–23 November 2018; pp. 1–15. [Google Scholar]
- Soleimani, M.; Kalhor, A.; Mirzadeh, H. Transformation-induced plasticity (TRIP) in advanced steels: A review. Mater. Sci. Eng. A
**2020**, 795, 140023. [Google Scholar] [CrossRef] - Sugimoto, K. Recent progress of low and medium carbon advanced martensitic steels. Metals
**2021**, 11, 652. [Google Scholar] [CrossRef] - Zackay, V.F.; Parker, E.R.; Fahr, D.; Bush, R. The enhancement of ductility in high-strength steels. Trans. Am. Soc. Met.
**1967**, 60, 252–259. [Google Scholar] - Grässel, O.; Krüger, L.; Frommeyer, G.; Meyer, L.W. High strength Fe-Mn-(Al, Si) TRIP/TWIP steels development − properties − application. Int. J. Plast.
**2000**, 16, 1391–1409. [Google Scholar] [CrossRef] - Sugimoto, K.; Sakaguchi, J.; Iida, T.; Kashima, T. Stretch-flangeability of a high-strength TRIP type bainitic steel. ISIJ Int.
**2000**, 40, 920–926. [Google Scholar] [CrossRef] - Sugimoto, K.; Tsunezawa, M.; Hojo, T.; Ikeda, S. Ductility of 0.1-0.6C-1.5Si-1.5Mn ultra high-strength low-alloy TRIP-aided sheet steels with bainitic ferrite matrix. ISIJ Int.
**2004**, 44, 1608–1614. [Google Scholar] [CrossRef] - Sugimoto, K.; Hashimoto, S.; Ikeda, S. Ultrahigh-strength low-alloy TRIP-aided sheet steels with bainitic ferrite matrix. In Proceedings of the Advanced High-Strength Sheet Steels for Automotive Applications Proceedings, Winter Park, CO, USA, 6–9 June 2004; pp. 63–70. [Google Scholar]
- Sugimoto, K.; Murata, M.; Muramatsu, T.; Mukai, Y. Formability of C-Si-Mn-Al-Nb-Mo ultra high-strength TRIP-aided sheet steels. ISIJ Int.
**2007**, 47, 1357–1362. [Google Scholar] [CrossRef] [Green Version] - Hojo, T.; Sugimoto, K.; Mukai, Y.; Ikeda, S. Effects of aluminum on delayed fracture properties of ultrahigh strength low alloy TRIP-aided steels. ISIJ Int.
**2008**, 48, 824–829. [Google Scholar] [CrossRef] [Green Version] - Sugimoto, K.; Murata, M.; Song, S. Formability of Al-Nb bearing ultrahigh-strength TRIP-aided sheet steels with bainitic ferrite and/or martensite matrix. ISIJ Int.
**2010**, 50, 162–168. [Google Scholar] [CrossRef] [Green Version] - Sugimoto, K.; Kobayashi, J. Advanced ultrahigh strength low alloy TRIP-aided steel sheet with excellent cold formability. J. Jpn. Soc. Technol. Plast
**2013**, 54, 949–953. [Google Scholar] [CrossRef] - Weiβensteiner, I.; Suppan, C.; Hebesberger, T.; Winkelhofer, F.; Clemens, H.; Maier-Kiener, V. Effect of morphological differences on the cold formability of an isothermally heat-treated advanced high-strength steel. JOM
**2018**, 70, 1567–1575. [Google Scholar] [CrossRef] [Green Version] - Tang, S.; Lan, H.; Liu, Z.; Wang, G. Enhancement of balance in strength, ductility and stretch flangeability by two-step austempering in a 1000 MPa grade cold rolled bainitic steel. Metals
**2021**, 11, 96. [Google Scholar] [CrossRef] - Zhou, Y.; Hojo, T.; Koyama, M.; Ajito, S.; Akiyama, E. Synergistic effects of hydrogen and deformation temperature on mechanical properties of TRIP-aided bainitic ferrite steel. Mater. Sci. Eng. A
**2022**, 842, 143070. [Google Scholar] [CrossRef] - Speer, J.G.; Edmonds, D.V.; Rizzo, F.C.; Matlock, D.K. Partitioning of carbon from supersaturated plates of ferrite, with application to steel processing and fundamentals of the bainite transformation. Curr. Opin. Solid State Mater. Sci.
**2004**, 8, 219–237. [Google Scholar] [CrossRef] - Speer, J.G.; De Moor, E.; Findley, K.O.; Matlock, B.C.; De Cooman, B.C.; Edmonds, D.V. Analysis of microstructure evolution in quenching and partitioning automotive sheet steel. Metall. Mater. Trans. A
**2011**, 42, 3591–3601. [Google Scholar] [CrossRef] - De Moor, E.; Speer, J.G.; Matlock, D.K.; Föjer, C.; Penning, J. Effect of Si, Al and Mo alloying on tensile properties obtained by quenching and partitioning. In Proceedings of the Materials Science & Technology Conference and Exhibition 2009 (MS&T’09), Pittsburgh, PA, USA, 25–29 October 2009; pp. 1554–1563. [Google Scholar]
- De Moor, E.; Speer, J.G.; Matlock, D.K. Effect of retained austenite on tensile behavior of AHSS revisited. In Proceedings of the Materials Science & Technology Conference and Exhibition 2011 (MS&T’11), Columbus, OH, USA, 16−20 October 2011; pp. 568–579. [Google Scholar]
- Seo, E.; Cho, L.; Estrin, Y.; De Cooman, B.C. Microstructure-mechanical properties relationships for quenching and partitioning (Q&P) processed steel. Acta Mater.
**2016**, 113, 124–139. [Google Scholar] - Kantanen, P.K.; Javaheri, V.; Somani, M.C.; Porter, D.A.; Komi, J.I. Effect of deformation and grain size on austenite decomposition during quenching and partitioning of (high) silicon-aluminum steels. Mater. Charact.
**2021**, 171, 110793. [Google Scholar] [CrossRef] - Celada-Casero, C.; Vercruysse, F.; Linde, B.; Smith, A.; Kok, P.; Sietsma, J.; Santofimia, M.J. Analysis of work hardening mechanisms in quenching and partitioning steels combining experiments with a 3D micro-mechanical model. Mater. Sci. Eng. A
**2022**, 846, 143301. [Google Scholar] [CrossRef] - Xia, P.; Vercruysse, F.; Celada-Casero, C.; Verleysen, P.; Petrov., R.H.; Sabirov, I.; Molina-Aldareguia, J.M.; Smith, A.; Linke, B.; Thiessen, R.; et al. Effect of alloying and microstructure on formability of advanced high-strength steels processed via quenching and partitioning. Mater. Sci. Eng. A
**2022**, 831, 142217. [Google Scholar] [CrossRef] - Garcia-Mateo, C.; Paul, G.; Somani, M.C.; Porter, D.A.; Bracke, L.; Latz, A.; De Andres, C.G.; Caballero, F.G. Transferring nanoscale bainite concept to lower contents: A prospective. Metals
**2017**, 7, 159. [Google Scholar] [CrossRef] [Green Version] - Zhu, K.; Mager, C.; Huang, M. Effect of substitution of Si by Al on the microstructure and mechanical properties of bainitic transformation-induced plasticity steels. J. Mater. Sci. Technol.
**2017**, 33, 1475–1486. [Google Scholar] [CrossRef] - Rana, R.; Chen, S.; Halder, S.; Das, S. Mechanical properties of a bainitic steel producible by hot rolling. Arch. Metall. Mater.
**2017**, 62, 2331–2338. [Google Scholar] [CrossRef] [Green Version] - Tian, J.; Xu, G.; Zhou, M.; Hu, H. Refined bainite microstructure and mechanical properties of a high-strength low-carbon bainitic steel. Steel Res. Int.
**2018**, 89, 1700469. [Google Scholar] [CrossRef] - He, S.H.; He, B.B.; Zhu, K.Y.; Huang, M.X. On the correlation among dislocation density, lath thickness and yield stress of bainite. Acta Mater.
**2017**, 135, 382–389. [Google Scholar] [CrossRef] - Tian, J.; Xu, G.; Zhou, M.; Hu, H.; Xue, Z. Effects of Al addition on bainite transformation and properties of high-strength carbide-free bainitic steels. J. Iron Steel Res. Int.
**2019**, 26, 846–855. [Google Scholar] [CrossRef] - Eres-Castellanos, A.; Caballero, F.G.; Garcis-Mateo, C. Stress or strain induced martensitic and bainitic transformations during ausforming process. Acta Mater.
**2020**, 189, 60–72. [Google Scholar] [CrossRef] - Lee, E.; Yu, C.; Lee, H.; Kim, J.; Suh, D. Influence of Si and Al contents and isothermal treatment condition on the microstructure and tensile properties in ultra-high strength Fe-0.2C-2.0Mn martensite-bainite complex phase steels. Kor. J. Met. Mater.
**2021**, 59, 602–612. [Google Scholar] [CrossRef] - Miller, R.I. Ultrafine-grained microstructures and mechanical properties of alloy steels. Metall. Trans.
**1972**, 3, 905–912. [Google Scholar] [CrossRef] - Furukawa, T. Dependence of strength-ductility characteristics on thermal history in low carbon, 5 wt-%Mn steels. Mater. Sci. Technol.
**1989**, 5, 465–470. [Google Scholar] [CrossRef] - Cao, W.; Wang, C.; Shi, J.; Wang, M.; Hui, W.; Dong, D. Microstructure and mechanical properties of Fe–0.2C–5Mn steel processed by ART-annealing. Mater. Sci. Eng. A
**2011**, 528, 6661–6666. [Google Scholar] [CrossRef] - Kang, S.; De Moor, E.; Speer, J.G. Retained austenite stabilization through solute partitioning during intercritical annealing in C-, Mn-, Al-, Si-, and Cr-alloyed steels. Metall. Mater. Trans. A
**2015**, 46, 1005–1011. [Google Scholar] [CrossRef] - Sugimoto, K.; Tanino, H.; Kobayashi, J. Warm stretch-formability of 0.2%C-1.5%Si-(1.5-5.0)%Mn TRIP-aided steels. Arch. Mater. Sci. Eng.
**2016**, 80, 5–15. [Google Scholar] [CrossRef] - Sugimoto, K.; Hidaka, S.; Tanino, H.; Kobayashi, J. Effects of Mn content on the warm stretch-flangeability of C-Si-Mn TRIP-aided steels. Steel Res. Int.
**2017**, 88, 1600482. [Google Scholar] [CrossRef] - Chang, Y.; Han, S.; Li, X.; Wang, C.; Zheng, G.; Dong, H. Effect of shearing clearance on formability of sheared edge of the third-generation automotive medium-Mn steel with metastable austenite. J. Mater. Process. Tech.
**2018**, 259, 216–227. [Google Scholar] [CrossRef] - Jing, S.; Ding, H.; Liu, M. Role of Al element in tailoring the austenite mechanical stability and tensile properties of medium Mn steels. J. Mater. Res. Technol.
**2022**, 20, 1414–1427. [Google Scholar] [CrossRef] - Cao, W.; Zhang, M.; Huang, C.; Xiao, S.; Dong, H.; Weng, Y. Ultrahigh Charpy impact toughness (−450J) achieved in high strength ferrite/martensite laminated steels. Sci. Rep.
**2017**, 7, 41459. [Google Scholar] [CrossRef] [Green Version] - Xu, Y.; Hu, Z.; Zou, Y.; Tan, X.; Han, D.; Chen, S.; Ma, D.; Misra, R.D.K. Effect of two-step intercritical annealing on microstructure and mechanical properties of hot-rolled medium manganese TRIP steel containing δ-ferrite. Mater. Sci. Eng. A
**2017**, 688, 40–55. [Google Scholar] [CrossRef] - Morawiec, M.; Ruiz-Jimenez, V.; Garcia-Mateo, C.; Jimenez, J.A.; Grajcar, A. Study of the isothermal bainitic transformation and austenite stability in an advanced Al-rich medium-Mn steel. Arch. Civil Mech. Eng.
**2022**, 22, 152. [Google Scholar] [CrossRef] - Kim, J.; Seo, E.; Kwon, M.; Kang, S.; De Cooman, B.C. Effect of quenching temperature on stretch flangeability of a medium Mn steel processed by quenching and partitioning. Mater. Sci. Eng. A
**2018**, 729, 276–284. [Google Scholar] [CrossRef] - Kaar, S.; Krizan, D.; Schneider, R.; Sommitsch, C. Impact of Si and Al on microstructural evolution and mechanical properties of lean medium Mn quenching and partitioning steel. Steel Res. Int.
**2020**, 91, 2000181. [Google Scholar] [CrossRef] - Kim, J.; Kwon, M.; Lee, J.; Lee, S.; Lee, K.; Suh, D. Influence of isothermal treatment prior to initial quenching of Q&P process on microstructure and mechanical properties of medium Mn steel. ISIJ Int.
**2021**, 60, 518–526. [Google Scholar] - Wallner, M.; Steineder, K.; Schneider, R.; Commenda, C.; Sommitsch, C. Effect of galvannealing on the microstructure and mechanical properties of a Si and Al alloyed medium-Mn quenching and partitioning steels. Mater. Sci. Eng. A
**2022**, 841, 143067. [Google Scholar] [CrossRef] - Gu, G.; Kim, J.; Lee, H.; Zargaran, A.; Koo, M.; Kim, S.; Lee, J.; Suh, D. Room temperature quenching (RT-Q&P) processed steel with chemically heterogeneous initial microstructure. Mater. Sci. Eng. A
**2022**, 851, 143651. [Google Scholar] - Kobayashi, J.; Pham, D.V.; Sugimoto, K. Stretch-flangeability of 1.5 GPa grade TRIP-aided martensitic cold rolled sheet steels. Steel Res. Int. (Special Edition: ICTP2011)
**2011**, 82, 598–603. [Google Scholar] - Sugimoto, K.; Kobayashi, J.; Pham, D.V. Advanced ultrahigh-strength TRIP-aided martensitic sheet steels for automotive applications. In Proceedings of the New Developments in Advanced High Strength Sheet Steels (AIST 2013), Vail, CO, USA, 23–27 June 2013; pp. 175–184. [Google Scholar]
- Pham, D.V.; Kobayashi, J.; Sugimoto, K. Effects of microalloying on stretch-flangeability of TRIP-aided martensitic sheet steel. ISIJ Int.
**2014**, 54, 1943–1951. [Google Scholar] [CrossRef] [Green Version] - Sugimoto, K.; Srivastava, A.K. Microstructure and mechanical properties of TRIP-aided martensitic steel. Metallogr. Microstruct. Anal.
**2015**, 4, 344–354. [Google Scholar] [CrossRef] [Green Version] - Kobayashi, J.; Nakashima, Y.; Sugimoto, K.; Itoh., G. Formabilities of C-Si-Al-Mn transformation-induced plasticity-aided martensitic sheet steel. Mater. Sci. Forum
**2016**, 838–839, 546–551. [Google Scholar] [CrossRef] - Jiang, H.; He, Y.; Lin, L.; Liu, R.; Zhang, Y.; Zheng, W.; Li, L. Microstructure and properties of auto-tempering ultra-high strength automotive steel under different thermal-processing conditions. Metals
**2021**, 11, 1121. [Google Scholar] [CrossRef] - Hanamura, T.; Torizuka, S.; Tamura, S.; Enokida, S.; Takechi, H. Effect of austenite grain size on transformation behavior, microstructure and mechanical properties of 0.1C-5Mn martensitic steel. ISIJ Int.
**2013**, 53, 2218–2225. [Google Scholar] [CrossRef] [Green Version] - He, B.; Liu, L.; Huang, M. Room-temperature quenching and partitioning steel. Metall. Mater. Trans. A
**2018**, 49A, 3167–3172. [Google Scholar] [CrossRef] - Sugimoto, K.; Tanino, H.; Kobayashi, J. Cold formabilities of martensite-type medium Mn steel. Metals
**2021**, 11, 1371. [Google Scholar] [CrossRef] - Seo, E.; Cho, L.; De Cooman, B.C. Application of quenching and partitioning (Q&P) processing to press hardening steel. Metall. Mater. Trans A
**2014**, 45A, 4022–4037. [Google Scholar] - Liang, J.; Zhao, Z.; Sun, B.; Lu, H.; Liang, J.; He, Q.; Chen, W.; Tang, D. A novel ultra-strong hot stamping steel treated by quenching and partitioning process. Mater. Sci. Technol.
**2018**, 34, 2241–2249. [Google Scholar] [CrossRef] - Wang, C.; Li, X.; Han, S.; Zhang, L.; Chang, Y.; Cao, W.; Dong, H. Warm stamping technology of the medium manganese steel. Steel Res. Int.
**2018**, 89, 1700360. [Google Scholar] [CrossRef] - Pan, H.; Cai, M.; Ding, H.; Sun, S.; Huang, H.; Zhang, Y. Ultrahigh strength-ductile medium-Mn steel auto-parts combining warm stamping and quenching & partitioning. Mater. Sci. Technol.
**2019**, 35, 807–814. [Google Scholar] - Ding, W.; Gong, Y.; Lu, Q.; Wang, J.; Wang, Z.; Li, W.; Jin, W. Improve bendability of a Cr-alloyed press-hardening steel through an in-line quenching and non-isothermal partitioning process. J. Manuf. Process.
**2022**, 84, 481–493. [Google Scholar] [CrossRef] - Echeverri, E.A.A.; Nishikawa, A.S.; Masoumi, M.; Pereira, H.B.; Marulanda, N.G.; Rossy, A.M.; Goldenstein, H.; Tschptschin, A.P. In site synchrotron X-ray diffraction and microstructural studies on cold and hot stamping combined with quenching & partitioning processing for development of third-generation advanced high strength steels. Metals
**2022**, 12, 174. [Google Scholar] - Chen, H.C.; Era, H.; Shimizu, M. Effect of phosphorus on the formation of retained austenite and mechanical properties in Si-containing low-carbon steel sheet. Metall. Trans. A
**1989**, 20A, 437–444. [Google Scholar] [CrossRef] - Barbé, L.; Verbeken, K.; Wettinck, E. Effect of the addition of P on the mechanical properties of low alloyed TRIP steel. ISIJ Int.
**2006**, 46, 1251–1267. [Google Scholar] [CrossRef] [Green Version] - Gao, J.; Ichikawa, M. Development of coated TRIP steels at Dofasco. In Proceedings of the Advanced High Strength Sheet Steels for Automotive Applications Proceedings, AIST, Winter Park, CO, USA, 6–9 June 2004; pp. 107–116. [Google Scholar]
- Sugimoto, K.; Mukherjee, M. Chapter 8: TRIP aided and complex phase steels. In Automotive Steels—Design, Metallurgy, Processing and Applications; Rana, R., Singh, S.B., Eds.; Woodhead Publishing: Cambridge, UK, 2016; pp. 107–116. [Google Scholar]
- Mahieu, J.; Van Dooren, D.; De Cooman, B.C. Influence of Si, Al and P on the thermodynamics and kinetics of the phase transformation related to continuous galvanization of TRIP-aided steels. In Proceedings of the 42nd Annual Conference of Metallurgists of CIM (COM2003), Vancouver, BC, Canada, 24–27 August 2003; pp. 21–35. [Google Scholar]
- Imai, N.; Komatsubara, N.; Kunishige, K. Effects of alloying elements and microstructure on the stability of retained austenite. CAMP-ISIJ
**1995**, 8, 76–80. [Google Scholar] - Imai, N.; Komatsubara, N.; Kunishige, K. Effects of alloying element and microstructure on mechanical properties of low-alloy TRIP steels. CAMP-ISIJ
**1995**, 8, 573–575. [Google Scholar] - De Meyer, M.; Vanderschueren, D.; De Cooman, B.C. The influence of the substitution of Si by Al on the properties of cold rolled C-Mn-Si TRIP steels. ISIJ Int.
**1999**, 39, 813–822. [Google Scholar] [CrossRef] - De, A.K.; Kircher, R.S.; Speer, J.G.; Matlock, D.K. Transformation behavior of retained austenite in TRIP steels as revealed by a specialized etching technique. In Proceedings of the Advanced High Strength Sheet Steels for Auto Applications Proceedings, Winter Park, CO, USA, 6–9 June 2004; pp. 338–347. [Google Scholar]
- Sugimoto, K.; Yu, B.; Mukai, Y.; Ikeda, S. Microstructure and formability of aluminum bearing TRIP-aided steels with annealed martensite matrix. ISIJ Int.
**2005**, 45, 1194–1200. [Google Scholar] [CrossRef] [Green Version] - Samek, L.; De Moor, F.; Penning, J.; De Cooman, B.C. Influence of alloying elements on the kinetics of strain-induced martensitic nucleation in low-alloy, multiphase high-strength steels. Metall. Mater. Trans. A
**2006**, 37A, 109–124. [Google Scholar] [CrossRef] - Lim, N.; Park, H.; Kim, S.; Park, C. Effects of aluminum on the microstructure and phase transformation of TRIP steels. Met. Mater. Int.
**2012**, 18, 647–654. [Google Scholar] [CrossRef] - Sugimoto, K.; Tsuruta, J.; Mukai, Y. The effects of cold-rolling strain on microstructure and formability of Al bearing TRIP-aided cold-rolled steel sheets with annealed bainite lath structure matrix. Tetsu Hagane
**2010**, 96, 29–36. [Google Scholar] [CrossRef] [Green Version] - Nagasaka, A.; Sugimoto, K.; Kobayashi, M.; Kobayashi, Y.; Hashimoto, S. Effests of structure and stability of retained austenite on deep drawability of TRIP-aided dual-phase sheet steels. Tetsu Hagane
**1999**, 85, 885–890. [Google Scholar] [CrossRef] [Green Version] - Takahashi, M. Development of high strength steels for automobiles. Nippon. Steel Tech. Rep.
**2003**, 88, 2–7. [Google Scholar] - Sugimoto, K.; Hojo, T.; Srivastava, A.K. Low and medium carbon advanced high-strength forging steels for automotive applications. Metals
**2019**, 9, 1263. [Google Scholar] [CrossRef] [Green Version] - Maruyama, H. X-ray measurement of retained austenite. Jpn. Soc. Heat Treat.
**1977**, 17, 198–204. [Google Scholar] - Dyson, D.J.; Holmes, B. Effect of alloying additions on the lattice parameter of austenite. J. Iron Steel Inst.
**1970**, 208, 469–474. [Google Scholar] - Koistinen, D.P.; Marburger, R.E. A general equation describing the extent of the austenite-martensite transformation in pure iron-carbon alloys and plain carbon steel. Acta Metall.
**1959**, 7, 59–60. [Google Scholar] [CrossRef] - Tamura, I. Strength of Steel; Nikkan-Kogyo Shinbunn Press: Tokyo, Japan, 1969; pp. 51–54. [Google Scholar]
- Sugimoto, K.; Kobayashi, M.; Hashimoto, S. Ductility and strain-induced transformation in a high-strength transformation-induced plasticity-aided dual-phase steel. Metall. Trans. A
**1992**, 23A, 3085–3091. [Google Scholar] [CrossRef] - Sherif, M.Y.; Mateo, C.G.; Sourmail, T.; Bhadeshia, H.K.D.H. Stability of retained austenite in TRIP-assisted steels. Mater. Sci. Technol.
**2004**, 20, 319–322. [Google Scholar] [CrossRef] [Green Version] - Sugimoto, K.; Kobayashi, J. Retained austenite characteristics and mechanical properties of C-Si/Al-Mn-Cr-B TRIP-aided martensitic steels. Unpublished work.
- Sugimoto, K.; Kobayashi, M.; Nagasaka, A.; Hashimoto, S. Warm stretch-formability of TRIP-aided dual-phase steel. ISIJ Int.
**1995**, 35, 1407–1414. [Google Scholar] [CrossRef] [Green Version] - Ehrhardt, B.; Gerber, T.; Schaumann, T.W. Approaches to microstructural design of TRIP and TRIP-aided cold rolled high strength steels. In Proceedings of the Advanced High Strength Sheet Steels for Automotive Applications (AHSSS) Proceedings, Winter Park, CO, USA, 6–9 June 2004; pp. 39–50. [Google Scholar]
- Alza, V.A.; Chavez, V.P. TRIP steels: Factors influencing their formation, mechanical properties and microstructure- A review. IOSR J. Mech. Civil Eng.
**2022**, 19, 37–60. [Google Scholar] - Dumay, A.; Chateau, J.; Allain, S.; Migot, S.; Bouuaziz, O. Influence of addition elements on the stacking-fault energy and mechanical properties of an austenitic Fe-Mn-C steel. Mater. Sci. Eng. A
**2008**, 483–484, 184–187. [Google Scholar] [CrossRef] - Saeed-Akbari, A.; Schwedt, A.; Bleck, W. Low stacking fault energy steels in the context of manganese-rich iron-based alloys. Scr. Mater.
**2012**, 66, 1024–1029. [Google Scholar] [CrossRef] - Leslie, W.C.; Rauch, G.C. Precipitation of carbides in low-carbon Fe-Al-C alloys. Metall. Trans. A
**1978**, 9, 343–349. [Google Scholar] [CrossRef] - Kinoshita, N.; Mura, T. Elastic fields of inclusions in anisotropic media. Phys. Status Solidi A
**1971**, 5, 759–768. [Google Scholar] [CrossRef] - Sakaki, T.; Sugimoto, K.; Fukuzato, T. Role of internal stress for continuous yielding of dual-phase steels. Acta Metall.
**1983**, 31, 1737–1746. [Google Scholar] [CrossRef] - Allen, N.P. In Iron and its Dilute Solid Solution; Spencer, C.W., Werner, F.E., Eds.; Willey Interscience: New York, NY, USA, 1963; pp. 271–308. [Google Scholar]
- Nagasaka, A.; Sugimoto, K.; Kobayashi, M.; Hashimoto, S. Effects of warm forming on stretch-flangeability of a TRIP-aided dual-phase sheet steel. Tetsu Hagane
**1997**, 83, 335–340. [Google Scholar] [CrossRef] [PubMed] [Green Version] - Sugimoto, K.; Kanda, A.; Kikuchi, R.; Hashimoto, S.; Kashima, T.; Ikeda, S. Ductility and formability of newly developed high strength low alloy TRIP-aided sheet steels with annealed martensite matrix. ISIJ Int.
**2002**, 42, 910–915. [Google Scholar] [CrossRef] - Liu, Z.; Hou, Y.; He, R.; Ye, Y.; Niu, C.; Min, J. Machine learning for extending capability of mechanical characterization to improve springback prediction of a quenching and partitioning steel. J. Mater. Process. Tech.
**2022**, 308, 117737. [Google Scholar] [CrossRef] - Gao, G.; Zhang, B.; Cheng, C.; Zhao, P.; Zhang, H.; Bai, B. Very high cycle fatigue behaviors of bainite/martensite multiphase steel treated by quenching-partitioning-tempering process. Int. J. Fatigue
**2016**, 92, 203–210. [Google Scholar] [CrossRef] - Sugimoto, K.; Hojo, T.; Srivastava, A.K. An overview of fatigue strength of case-hardening TRIP-aided martensitic steels. Metals
**2018**, 8, 355. [Google Scholar] [CrossRef] [Green Version] - Zhou, Q.; Qian, L.; Meng, J.; Zhao, L. The fatigue properties, microstructural evolution and crack behaviors of low-carbon carbide-free bainitic steel during low-cycle fatigue. Mater. Sci. Eng. A
**2021**, 820, 141571. [Google Scholar] [CrossRef] - Paravicini Bagliani, E.; Santofimia, M.J.; Zhao, L.; Sietsma, J.; Anelli, E. Microstructure, tensile and toughness properties after quenching and partitioning treatments of a medium-carbon steel. Mater. Sci. Eng. A
**2013**, 559, 486–495. [Google Scholar] [CrossRef] - Somani, M.C.; Porter, D.A.; Karjalainen, L.P.; Suikkanen, P.P.; Misra, R.D.K. Process design for tough ductile martensitic steels through direct quenching and partitioning. Matter. Today: Proceedings
**2015**, 2S, S631–S634. [Google Scholar] [CrossRef] - Frómeta, D.; Parareda, S.; Lara, A.; Molas, S.; Casellas, D.; Jonsén, P.; Calvo, J. Identification of fracture toughness parameters to understand the fracture resistance of advanced high strength sheet steels. Eng. Fract. Mech.
**2020**, 229, 106949. [Google Scholar] [CrossRef] - Zhang, Y.; Hui, W.; Zhao, X.; Wang, C.; Cao, W.; Dong, H. Effect of reverted austenite fraction on hydrogen embrittlement of TRIP-aided medium Mn steel (0.1C-5Mn). Eng. Fail. Anal.
**2019**, 97, 605–616. [Google Scholar] [CrossRef] - Hojo, T.; Akiyama, E.; Saito, H.; Shiro, A.; Yasuda, R.; Shobu, T.; Kinugasa, J.; Yuse, F. Effects of residual stress and plastic strain on hydrogen embrittlement of a stretch-formed TRIP-aided martensitic steel sheet. Corr. Sci.
**2020**, 177, 108957. [Google Scholar] [CrossRef] - Li, W.; Ma, L.; Peng, P.; Jia, Q.; Wan, Z.; Zhu, Y.; Guo, W. Microstructural evolution and deformation behavior of fiber laser welded QP980 steel joint. Mater. Sci. Eng. A
**2018**, 717, 124–133. [Google Scholar] [CrossRef] - Królicka, A.; Ambroziak, A.; Żak, A. Welding capabilities of nanostructured carbide-free bainite: Review of welding methods, materials, problems, and perspectives. Appl. Sci.
**2019**, 9, 3798. [Google Scholar] [CrossRef] [Green Version] - Cao, Y.; Zhao, L.; Peng, Y.; Song, L.; Zhong, M.; Ma, C.; Tian, Z. Microstructure and mechanical properties of simulated heat affected zone of laser welded medium-Mn steel. ISIJ Int.
**2019**, 60, 2266–2275. [Google Scholar] [CrossRef] - Zhang, B.; Dong, Y.; Du, Y.; Misra, R.D.K.; Wu, H.; Wang, X.; Zhao, W.; Du, L. Microstructure and formability performance of fiber laser welded 1.2 GPa grade hot-rolled TRIP steel joints. Optics Laser Technol.
**2021**, 143, 107341. [Google Scholar] [CrossRef] - Tümer, M.; Schneider-Bröskamp, C.; Enzinger, N. Fusion welding of ultra-high strength structural steel—A review. J. Manuf. Process
**2022**, 82, 203–229. [Google Scholar] [CrossRef] - Wu, X.; Lin, H.; Wang, Y.; Jiang, H. Hydrogen embrittlement and fracture mechanism of friction stir welded quenching and partitioning 980 steel. Mater. Sci. Eng. A
**2021**, 802, 140683. [Google Scholar] [CrossRef] - Ling, Z.; Chen, T.; Wang, M.; Kong, L. Reducing liquid metal embrittlement cracking in resistance spot welding of Q&P980 steel. Mater. Manuf. Process.
**2020**, 35, 1392–1399. [Google Scholar] - Stadler, M.; Gruber, M.; Schnitzer, R.; Hofer, C. Microstructural characterization of a double pulse resistance spot welded 1200 MPa TBF steel. Weld. World
**2020**, 64, 335–343. [Google Scholar] - Wang, Z.; Liu, M.; Zhang, H.; Xie, G.; Xue, P.; Wu, L.; Zhang, Z.; Ni, D.; Xiao, B.; Ma, Z. Welding behavior of an ultrahigh-strength quenching and partitioning steel by fusion and solid-state welding methods. J. Mater. Res. Technol.
**2022**, 17, 1289–1301. [Google Scholar] [CrossRef] - Nagasaka, A.; Hojo, T.; Aoki, K.; Koyama, H.; Shimizu, A. Spot-welded tensile properties in automobile ultrahigh-strength TRIP-aided martensitic steel sheet. ISIJ Int.
**2021**, 61, 599–607. [Google Scholar] [CrossRef] - Raedt, H.; Speckenheuer, U.; Vollrath, K. New forged steels, Energy-efficient solutions for stronger parts. Automobiltech Z. (ATZ)
**2012**, 114, 5–9. [Google Scholar] - Raedt, H.W.; Wilke, F.; Ernst, C.S. Lightweight forging initiative—Phase II: Lightweight design potential in a light commercial vehicle. Automobiltech Z. (ATZ)
**2016**, 118, 48–52. [Google Scholar] [CrossRef] - Buchmayr, B. Critical assessment 22: Bainitic forging steels. Mater. Sci. Technol.
**2016**, 32, 517–522. [Google Scholar] [CrossRef] [Green Version] - Raedt, H.W.; Wurm, T.; Busse, A. The lightweight forging initiative—Phase III: Lightweight forging design for hybrid cars and heavy-duty trucks. Automobiltech Z. (ATZ)
**2019**, 120, 54–59. [Google Scholar] [CrossRef] - Sugimoto, K.; Sato, S.; Kobayashi, J.; Srivastava, A.K. Effects of Cr and Mo on mechanical properties of hot-forged medium carbon TRIP-aided bainitic ferrite steels. Metals
**2019**, 9, 1066. [Google Scholar] [CrossRef] [Green Version] - Gramlich, A.; Schmiedl, T.; Schönborn, S.; Melz, T.; Bleck, W. Development of air-hardening martensitic steels. Mater. Sci. Eng. A
**2020**, 784, 139321. [Google Scholar] [CrossRef]

**Figure 2.**Heat treatment diagrams of (

**a**) TBF, CFB and TM steels [52] and (

**b**) one-step and two-step Q&P steels [20]. T

_{IT}, T

_{Q}, T

_{P}, Ac

_{3}, M

_{s}, and M

_{f}are isothermal transformation temperature, quenching temperature, partitioning temperature, the austenite-finish temperature on heating, martensite-start temperature, and martensite-finish temperatures, respectively.

**Figure 3.**Illustration of the typical microstructure of various third-generation AHSSs [81]. (

**a**): TBF, CFB, and BF-MMn steels (T

_{IT}> M

_{s}); (

**b**): TBF, CFB, one-step Q&P, and BF-MMn steels (T

_{IT}= M

_{s}− M

_{f}), and two-step Q&P and two-step Q&P-MMn steels (T

_{P}> M

_{s}); (

**c**): TM and M-MMn steels (T

_{IT}< M

_{f}). α

_{bf}, α

_{m}, α

_{m}*, γ

_{R}, θ, and MA represent bainitic ferrite, primary coarse soft martensite, secondary fine hard martensite, retained austenite, carbide, and MA phase (a mixture of α

_{m}* and film-like γ

_{R}), respectively.

**Figure 4.**Variations in initial volume fraction (fγ

_{0}, ●) and carbon concentration (Cγ

_{0}, ○) of retained austenite as a function of isothermal transformation temperature (T

_{IT}) in 0.20C-1.59Si-1.50Mn-0.05Nb (mass%) TBF and TM steels [51]. The holding time of the IT process is 1000 s. This figure is reproduced based on reference [51].

**Figure 5.**(

**a**) Illustration of variations in volume fractions of different phases as a function of quenching temperature (T

_{Q}) in two-step Q&P steel. (

**b**) Variations in initial retained austenite fraction with T

_{Q}in 0.2C-4.0Mn-1.6Si-1.0Cr (M

_{s}= 273 °C) and 0.3C-1.0Mn-1.6Si-1.0Cr (M

_{s}= 235 °C) two-step Q&P steels [23]. (

**a**) is reproduced based on references [19,20], in which fα

_{m}and fγ are volume fractions of primary coarse soft martensite and austenite as functions of T

_{Q}prior to partitioning. The final or initial austenite fraction (fγ

_{0}) at room temperature is given by a red, bold, solid line. fα

_{m}’ = fα

_{bf}(bainitic ferrite fraction) + fα

_{m}* (secondary fine hard martensite fraction). (

**b**) is reprinted with permission from Elsevier, copyright 2022.

**Figure 6.**Relationships between strain-induced transformation factor (k) and initial carbon concentration of retained austenite (Cγ

_{0}) in 0.2C-1.5Si-1.2Mn-0.2Cr-(0.022-1.22)Al (0Al, 0.7Al, 1.2Al) TM steels (●) [55,88], 0.2C-1.5Si-(1.5-5.0)Mn (1.5Mn, 3Mn, and 5Mn) M-MMn (■) [59] and D-MMn (□) [39,40] steels, 0.2C-1.5Si-1.5Mn-(0-1.0)Cr-(0-0.2)Mo (Cr-Mo) TM steels (▲) [53] and 0.2C-1.5Si-1.5Mn (1.5Si) TBF steel (△) [51], and 0.2C-1.5Si-1.5Mn TPF (⊠) [89] and TAM (⊠) [89] steels.

**Figure 7.**(

**a**) Effects of microalloying elements on the time-temperature transformation curve of steel [90]. (

**b**) Calculated T

_{0}curves and measured carbon concentration of retained austenite (Cγ

_{0}) on 19C-1.54Si-1.51Mn-0.04Al (0Al), 0.5Al: 0.20C-0.99Si-1.51Mn-0.49Al (0.5Al), and 0.20C-0.49Si-1.50Mn-0.99Al (1.0Al) TPF steels [75]. (

**a**) is reprinted with permission from AIST, copyright 2022. (

**b**) is reproduced with permission from ISIJ, copyright 2022.

**Figure 8.**Quenching temperature dependences of yield stress (YS), tensile strength (UTS), and total elongation (TE) of 0.2C-4.0Mn-1.6Si-1.0Cr (0.2C) and 0.3C-4.0Mn-1.6Si-1.0Cr (0.3C) two-step Q&P steels partitioned at T

_{P}= 450 °C for 300 s after quenching [23]. M

_{s}s of the 0.2C and 0.3C Q&P steels are 273 and 235 °C, respectively. This figure is reprinted with permission from Elsevier, copyright 2022.

**Figure 9.**Variations in (

**a**) yield stress (YS), tensile strength (TS), (

**b**) uniform elongation (UEl), and total elongation (TEl); (

**c**) product of TS and TEl (TS×TEl) as a function of isothermal transformation temperature (T

_{IT}) in 0.20C-1.50Si-1.50Mn-0.05Nb TBF and TM steels. This figure is reproduced based on reference [51].

**Figure 10.**(

**a**) Engineering stress–strain (σ-ε) curves of 0.20C-1.50Si-1.24Mn-0.20Cr-0.022Al-0.0028 (0Al), 0.20C-0.73Si-1.24Mn-0.19Cr-0.72Al-0.0027B (0.7Al), and 0.20C-0.20Si-1.24Mn-0.20Cr-1.22Al-0.0026B (1.2Al) TM steels [55,88]. (

**b**) Combination of the tensile strength (TS) and total elongation (TEl) of 0Al, 0.7Al, and 1.2Al TM steels (●) [55,88], 0.2C-1.5Si-1.5Mn-(0-1.0)Cr-(0-0.2)Mo (Cr-Mo) TM steels (▲) [53], 0.2C-1.5Si-(1.5-5.0)Mn (1.5Mn, 3Mn, and 5Mn) M-MMn (■) [59] and D-MMn (□) [39,40] steels, 0.2C-1.5Mn-0.99Si-0.49Al (0.5Al) TBF steel (○), 0.2C-1.5Mn-1.00Si-0.48Al-0.049Nb (0.5Al-0.05Nb) TBF steel (▽) [14], and 0.2C-1.5Si-1.5Mn-0.05Nb (1.5Si) TBF steel (△) [51]. (

**b**) is produced based on references [14,39,40,51,53,55,59,88].

**Figure 11.**(

**a**) Relationship between tensile strength (TS) and initial retained austenite fraction (fγ

_{0}). (

**b**) Relationships between product of TS and total elongation (TS×TEl) and strain-induced transformation factor (k) in 0.20C-1.50Si-1.24Mn-0.20Cr-0.022Al-0.0028 (0Al), 0.20C-0.73Si-1.24Mn-0.19Cr-0.72Al-9.0027B (0.7Al), and 0.20C-0.20Si-1.24Mn-0.20Cr-1.22Al-0.0026B (1.2Al) TM steels (●) [55,88], 0.2C-1.5Si-1.5Mn-(0-1.0)Cr-(0-0.2)Mo (Cr-Mo) TM steels (▲) [53], 0.2C-1.5Si-(1.5-5.0)Mn (1.5Mn, 3Mn and 5Mn) M-MMn (■) [59] and D-MMn (□) [39,40] steels, 0.2C-1.5Mn-0.99Si-0.49Al (0.5Al) TBF steel (○), 0.2C-1.5Mn-1.00Si-0.48Al-0.049Nb (0.5Al-0.05Nb) TBF steel (▽) [14], and 0.2C-1.5Si-1.5Mn-0.05Nb (1.5Si) TBF steel (△) [51]. (

**a**,

**b**) are produced from the results in references [14,39,40,51,53,55,59,88].

**Figure 12.**(

**a**) Isothermal transformation (T

_{IT}) and partitioning temperature (T

_{P}) dependences in maximum stretch height (H

_{max}) in 0.20C-1.50Si-1.50Mn-0.05Nb TBF and TM steels subjected to isothermal transformation (IT) process (●) and TM steel subjected to direct quenching and then partitioning (DQ&P) process (◆) [51]. (

**b**) Relationship between H

_{max}and tensile strength (TS) in the IT-processed TBF and TM steels (●) [51], DQ&P-processed TM steel (◆) [51], 0.2C-1.5Si-(1.5-5.0)Mn (1.5Mn, 3Mn, and 5Mn) M-MMn (■) [59] and D-MMn (□) [39] steels, 0.082C-0.88Si-2.0Mn ferrite-martensite DP steel (◇) [14,51,52], and 0.23C-0.19Si-1.29Mn-0.21Cr-0.003B 22MnB5 Q&T steel (⊠) [14,52]. (

**a**) is reproduced based on reference [51]. (

**b**) is reproduced based on references [14,39,51,52,59].

**Figure 13.**(

**a**) Maximum stretch height–tensile strength

**(**H

_{max}−TS) relation in 0.20C-1.50Si-1.24Mn-0.20Cr-0.022Al-0.0028B (0Al), 0.20C-0.73Si-1.24Mn-0.19Cr-0.72Al-9.0027B (0.7Al), and 0.20C-0.20Si-1.24Mn-0.20Cr-1.22Al-0.0026B (1.2Al) TM steels (●) [55,88], 0.2C-1.5Si-1.5Mn-(0-1.0)Cr-(0-0.2)Mo (Cr-Mo) TM steels (▲) [53], and 0.2C-1.5Si-1.5Mn-0.05Nb (1.5Si) TBF steel (△) [51]. (

**b**) Relationship between H

_{max}and total elongation (TEl) in these steels and 0.2C-1.5Si-(1.5-5.0)Mn (1.5Mn, 3Mn and 5Mn) M-MMn (■) [59] and D-MMn (□) [39] steels. (

**a**,

**b**) are produced based on references [39,51,53,55,59,88].

**Figure 14.**(

**a**) Isothermal transformation and partitioning temperature (T

_{IT}and T

_{P}) dependences of hole-expansion ratio (HER) of 0.20C-1.50Si-1.50Mn-0.05Nb TBF and TM steels subjected to isothermal transformation (IT) process (●) and TM steel subjected to a direct quenching and partitioning (DQ&P) process (◆) [51]. (

**b**) HER−TS (tensile strength) relation in the IT-processed TBF and TM steels (●) [51], the DQ&P-processed TM steel (◆) [51], 0.2C-1.5Si-(1.5-5.0)Mn (1.5Mn, 3Mn, and 5Mn) M-MMn (■) [59] and D-MMn (□) [40] steels, 0.082C-0.88Si-2.0Mn ferrite-martensite DP steel (◇) [14,51,52], and 0.23C-0.19Si-1.29Mn-0.21Cr-0.003B 22MnB5 Q&T steel (⊠) [14,52]. (

**a**) is reproduced based on reference [51]. (

**b**) is produced based on references [14,40,51,52,59].

**Figure 15.**(

**a**) Relationship between hole-expansion ratio (HER) and tensile strength (TS) in 0.20C-1.50Si-1.24Mn-0.20Cr-0.022Al-0.0028 (0Al), 0.20C-0.73Si-1.24Mn-0.19Cr-0.72Al-9.0027B (0.7Al), and 0.20C-0.20Si-1.24Mn-0.20Cr-1.22Al-0.0026B (1.2Al) TM steels (●) [55,88], 0.2C-1.5Si-1.5Mn-(0-1.0)Cr-(0-0.2)Mo (Cr-Mo) TM steels (▲) [53], 0.20C-0.99Si-1.51Mn-0.49Al (0.5Al) TBF steel (○) [14], 0.20C-1.00Si-1.50Mn-0.48Al-0.049Nb (0.5Al-0.05Nb) TBF steel (▽) [14], and 0.2C-1.5Si-1.5Mn-0.05Nb (1.5Si) TBF steel (△) [51]. (

**b**) Variations in the product of TS and HER (TS×HER) as a function of isothermal transformation temperature (T

_{IT}) in 0.19C-1.54Si-1.51Mn-0.04Al (0Al), 0.5Al, 0.21C-1.50Si-1.51Mn-0.04Al-0.048Nb (0Al-0.05Nb), and 0.5Al-0.05Nb TBF steels [14]. (

**a**) is produced based on references [14,51,53,55,88]. (

**b**) is reproduced with permission from ISIJ, copyright 2022.

**Figure 16.**(

**a**) Relationship between hole expansion ratio (HER) and a ratio of shear section length to sheet thickness (ss/t) and (

**b**) relationship between ss/t and reduction of area (RA) in 0.20C-1.50Si-1.24Mn-0.20Cr-0.022Al-0.0028 (0Al), 0.20C-0.73Si-1.24Mn-0.19Cr-0.72Al-9.0027B (0.7Al), and 0.20C-0.20Si-1.24Mn-0.20Cr-1.22Al-0.0026B (1.2Al) TM steels (●) [55,88], 0.2C-1.5Si-1.5Mn-(0-1.0)Cr-(0-0.2)Mo (Cr-Mo) TM steels (▲) [53], 0.2C-1.5Si-(1.5-5.0)Mn (1.5Mn, 3Mn, and 5Mn) M-MMn steels (■) [59], 0.20C-1.50Si-1.50Mn-0.05Nb (1.5Si) TBF steel (△) [51], and 0.2C-1.00Si-0.48Al-0.049Nb (0.5Al-0.05Nb) TBF steels (▽) [14]. (

**a**,

**b**) are produced based on references [51,53,55,59,88].

**Figure 17.**Variations in (

**a**) average bending angle, (

**b**) yield stress (YS), ultimate tensile strength (UTS), and total elongation (TE), and (

**c**) volume fractions of retained austenite and martensite as a function of austempering temperature in 0.34C-1.65Si-1.94Mn-1.07Cr CFB steel [29]. M

_{s}of the steel is 329 °C.

**Figure 18.**Relationships between minimum bending radius (R

_{min}) and tensile strength (TS) in (

**a**) 0.21C-1.449Si-1.50Mn-1.0Cr-0.05Nb TBF and TM steels (●), 0.082C-0.88Si-2.0Mn ferrite-martensite DP steel (◇), and 22MnB5 Q&T steel (⊠) [52] and (

**b**) (0.1-0.6)C-1.5Si-1.5Mn TPF (○), TAM (□), and TBF (△) steels [99]. (

**a**) is reprinted with permission from AIST, copyright 2022. (

**b**) is reprinted with permission from ISIJ, copyright 2022.

**Figure 19.**Variations in minimum bending radius (R

_{min}) as a function of isothermal transformation temperature (T

_{IT}) in (

**a**) TAM and (

**b**) TPF steels with the chemistry of 0.19C-1.54Si-1.51Mn-0.04Al (0Al), 0.20C-0.99Si-1.51Mn-0.49Al (0.5Al), and 0.20C-0.49Si-1.50Mn-0.99Al (1.0Al) [75]. (

**a**,

**b**) are reprinted with permission from ISIJ, copyright 2022.

**Table 1.**Chemical composition (in mass%), measured properties, and references for low-carbon Si/Al-Mn first- and third-generation AHSSs used in various kinds of research.

Gen. | Steel | Chemical Composition | Property | Ref. |
---|---|---|---|---|

1st Gen. | TPF | 0.25C-1.28Si-1.67Mn-0.03Al, 0.18C-0.02Si-1.56Mn-1.73Al | 1 | [70] |

0.21C-2.10Si-1.52Mn-0.022Al, 0.22C-0.01Si-1.49Mn-2.02Al | 1, 2 | [71] | ||

0.21C-2.10Si-1.52Mn-0.022Al, 0.22C-0.01Si-1.49Mn-2.02Al | 1, 2, 4 | [72] | ||

0.19C-1.46Si-1.57Mn-0.06Al, 0.31C-0.34Si-1.57Mn-1.23Al | 1, 2 | [73] | ||

(0.14-0.21)C-(0.34-1.47)Si-1.5Mn-(0.03-0.99)Al | 1, 2 | [74] | ||

0.20C-(0.49-1.54)Si-1.5Mn-(0.04-0.99)Al | 1, 2, 4, 5 | [75] | ||

(0.19-0.25)C-(0.09-1.45)Si-1.7Mn-(0.03-1.49)Al | 1 | [76] | ||

0.20C-1.87Si-1.99Mn-(0.04-2.0)Al | 1 | [77] | ||

0.20C-(0.49-1.50)Si-1.5Mn-(0.04-0.99)Al | 1, 2, 4 | [78] | ||

TAM | 0.20C-(0.49-1.54)Si-1.5Mn-(0.04-0.99)Al | 1, 2, 4, 5 | [75] | |

0.20C-(0.48-1.50)Si-1.5Mn-(0.04-0.99)Al | 1, 2, 4 | [78] | ||

3rd Gen. | TBF | 0.20C-(0.49-1.54)Si-(1.48-1.51)Mn-(0.04-0.99)Al-(0-0.05)Nb-(0-0.20)Mo | 1, 2, 4 | [12] |

0.20C-(0.49-1.51)Si-(1.51-2.51)Mn-(0.04-0.99)Al | 1, 2 | [13] | ||

0.20C-(0.99-1.54)Si-1.5Mn-(0.04-0.49)Al-(0-0.05)Nb | 1, 2, 4 | [14] | ||

Q&P | 0.24C-1.45Si-1.61Mn-0.30Al, 0.25C-0.55Si-1.70Mn-0.69Al | 1,2 | [21] | |

0.24C-0.12Si-1.60Mn-1.41Al-0.17Mo | 2 | [22] | ||

0.30C-(0.48-0.99)Si-(1.86-2.00)Mn-(0.01-1.10)Al-(1.01-2.20)Cr | 1 | [24] | ||

CFB | 0.25C-(0.08-1.09)Si-2.07Mn-(0.021-1.54)Al | 1, 2 | [28] | |

0.25C-2.1Mn-(0.02-1.54)Al | 1, 2 | [31] | ||

0.22C-(1.79-1.82)Si-(1.98-2.04)Mn-(0-0.50)Al-1.0Cr-0.23Mo | 1, 2 | [32] | ||

0.2C-1.55Si-2.0Mn, 0.2C-0.77Si-2.0Mn-0.76Al | 1, 2 | [34] | ||

D-MMn | (0.1-0.3)C-(0-1.5)Si-(2-5)Mn-(0-1.5)Al-(0-1.5)Cr | 1 | [38] | |

(0.18-0.19)C-(7.66-7.93)Mn-(0-2.79)Al | 1, 2 | [42] | ||

BF-MMn | 0.18C-0.23Si-3.6Mn-1.7Al-0.2Mo-0.04Nb | 1 | [45] | |

Q&P-MMn | 0.173C-4.46Mn-1.47Si-0.03Al, 0.195C-4.52Mn-0.04Si-1.31Al | 1, 2 | [47] | |

0.2C-1.50Si-4.02Mn-0.02Al, 0.2C-0.08Si-4.04Mn-1.46Al | 1, 2 | [49] | ||

TM | 0.20C-(0.20-1.50)Si-1.24Mn-(0.02-1.22)Al-0.2Cr-(0.003-0.005) Ti-(0.003-0.005)B | 1, 2, 3, 4 | [55] | |

1: microstructure, 2: tensile properties, 3: stretch formability, 4: stretch-flangeability, 5: bendability. |

Publisher’s Note: MDPI stays neutral with regard to jurisdictional claims in published maps and institutional affiliations. |

© 2022 by the author. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (https://creativecommons.org/licenses/by/4.0/).

## Share and Cite

**MDPI and ACS Style**

Sugimoto, K.-i.
Effects of Partial Replacement of Si by Al on Cold Formability in Two Groups of Low-Carbon Third-Generation Advanced High-Strength Steel Sheet: A Review. *Metals* **2022**, *12*, 2069.
https://doi.org/10.3390/met12122069

**AMA Style**

Sugimoto K-i.
Effects of Partial Replacement of Si by Al on Cold Formability in Two Groups of Low-Carbon Third-Generation Advanced High-Strength Steel Sheet: A Review. *Metals*. 2022; 12(12):2069.
https://doi.org/10.3390/met12122069

**Chicago/Turabian Style**

Sugimoto, Koh-ichi.
2022. "Effects of Partial Replacement of Si by Al on Cold Formability in Two Groups of Low-Carbon Third-Generation Advanced High-Strength Steel Sheet: A Review" *Metals* 12, no. 12: 2069.
https://doi.org/10.3390/met12122069