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Review

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

by
Koh-ichi Sugimoto
School of Science and Technology, Shinshu University, Nagano 380-8553, Japan
Metals 2022, 12(12), 2069; https://doi.org/10.3390/met12122069
Submission received: 9 November 2022 / Revised: 25 November 2022 / Accepted: 28 November 2022 / Published: 1 December 2022
(This article belongs to the Special Issue Heat Treatment Process and Application of High-Strength Steel)
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Abstract

:
Partial replacement of Si by Al improves the coatability (or galvanizing property) of Si-Mn advanced high-strength steel (AHSS) sheets. In this paper, the effects of the partial replacement on the microstructure, tensile property, and cold formability are reported for the low-carbon third-generation AHSS sheets, which are classified into two groups, “Group I” and “Group II”. The partial replacement by 1.2 mass% Al increases the carbon concentration or mechanical stability of retained austenite and decreases its volume fraction in the AHSSs, compared to Al-free AHSSs. The partial replacement deteriorates the tensile ductility and stretch formability in the AHSSs with a tensile strength above 1.2 GPa. On the other hand, it achieves the same excellent stretch-flangeability as Al-free AHSSs. A complex addition of Al and Nb/Mo further enhances the stretch-flangeability. The cold formabilities are related to the heat treatment condition and microstructural and tensile properties, and the stress state.

1. Introduction

Since the 1970s, the first-, second-, and third-generation advanced high-strength steel (AHSS) sheets have been developed for weight reduction and enhancement of collision safety of automobiles [1,2,3,4,5,6]. In these AHSSs, the performance of the mechanical properties are principally enhanced by transformation-induced plasticity (TRIP) [7] and/or twinning-induced plasticity (TWIP) [8] of metastable retained austenite (γR), 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 (fγ0) (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.
The third-generation AHSSs are classified into “Group I” and “Group II” by a kind of matrix structure and TS level [1,3,4,5,6].
Group I: TRIP-aided bainitic ferrite (TBF) steel [9,10,11,12,13,14,15,16,17,18]; one-step and two-step quenching and partitioning (Q&P) steels [19,20,21,22,23,24,25,26]; carbide-free bainitic (CFB) steel [27,28,29,30,31,32,33,34]; and duplex type [35,36,37,38,39,40,41,42], laminate type [43,44], bainitic ferrite-type [45], and Q&P-type [46,47,48,49,50] medium-manganese (D-MMn, L-MMn, BF-MMn, and Q&P-MMn) steels.
Group II: TRIP-aided martensitic (TM) steel [15,51,52,53,54,55,56] and martensite-type medium-manganese (M-MMn) steel [50,57,58,59].
In Group I, the kind of matrix structure is bainitic ferrite (or bainitic ferrite/martensite mixture) and the TS level is higher than 1.0 GPa, except for D–MMn and L-MMn steels with the duplex structure of an annealed martensite and reverted austenite [35,36,37,38,39,40,41,42] and laminate structure of the δ-ferrite and complex structure of α-ferrite plus reverted austenite [43,44], respectively. On the other hand, the main matrix structure of Group II is martensite, and its TS level is higher than 1.5 GPa [15,51,52,53,54,55,56,57,58,59]. In Group II, martensite–austenite (MA) constituent or phase plays an important role in the strain-hardening behavior [52]. The ultimate goals of both group sheets of steel are to achieve the TS×TEl above 30 GPa% and excellent cold- and warm formability. Due to its excellent nature, the heat treatment process of the third-generation AHSSs is also applied to hot and warm stamping products [60,61,62,63,64,65].
Metals 12 02069 g001 550
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 (fγ0) 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].
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 (fγ0) 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].
Metals 12 02069 g001
The TS×TEl of the AHSSs is controlled by various alloying elements such as C, Si, Al, Mn, Cr, Mo, Ni, B, and P [1,5,11,12,13,14,21,22,24,26,28,31,32,33,34,35,36,38,39,42,43,44,45,46,47,48,49,53,54,55,56,59], and the heat-treatment process [1,4,9,10,11,12,13,14,15,16,17,19,20,22,23,24,25,26,27,28,29,30,31,32,33,34,35,36,37,38,39,42,43,44,45,46,47,48,49,50,51,52,53,54,56,57,58,59]. In this case, low-carbon content below 0.25 mass% is preferred to keep high weldability. The addition of Si and/or Al suppresses the carbide formation and resultantly increases the volume fraction of metastable γR 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].
To promote the application of galvanized AHSS sheets to automotive parts, many researchers investigate the effect of partial replacement of Si by Al on the microstructural and mechanical properties, such as tensile property and formability in low-carbon TBF [12,13,14], Q&P [21,22,24], CFB [28,31,32,34], D-MMn [38,42], BF-MMn [45], Q&P-MMn [47,49], and TM [55] steels, in the same manner as the first-generation AHSSs, such as TRIP-aided polygonal ferrite (TPF) [70,71,72,73,74,75,76,77,78] and TRIP-aided annealed martensite (TAM) steels [75,78] (Table 1). Unfortunately, most of the mechanical properties are focused on tensile properties, not formability. In this paper, the influences of the partial replacement of Si by Al on the cold formability, such as stretch formability, stretch-flangeability, and bendability of two groups of low-carbon third-generation AHSSs, are summarized, along with the microstructural and tensile properties. For ease of understanding, these characteristics are stated separately for low-carbon “Si-Mn” and “Si/Al-Mn” third-generation AHSSs. Unfortunately, there is not any research on deep drawability in the third-generation AHSS, except for the first-generation AHSSs such as TPF [79,80] and TAM [79] steels. Thus, the deep drawability of the third-generation AHSSs is omitted in this review.

2. Microstructure and Retained Austenite Characteristics

2.1. C-Si-Mn steel

In general, microstructure and retained austenite properties of the third-generation AHSSs are strongly controlled by the chemical composition and following heat treatment conditions:
Representative heat treatment diagrams of the third-generation AHSSs are shown in Figure 2. Low-carbon TBF and TM steels can be produced by austenitizing and subsequent IT processes. For the TBF steel, IT processes at temperatures (TIT) above Ms (IT process (i)) or between Ms and Mf (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 Mf (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].
The major phase of the TBF steel is bainitic ferrite (αbf) for the IT process (i) (Figure 3a) [81], and it a mixture of αbf and primary coarse soft martensite (αm) for the IT process (ii) (Figure 3b) [81]. In the TBF steel, an initial γR fraction (fγ0) increases with increasing TIT (Figure 4) [51]. The highest initial carbon concentration (Cγ0) is obtained in the TBF steel subjected to the IT process at the temperatures between Ms and Mf. In the TBF steel subjected to the IT process (ii), a small amount of MA phase (a mixture of secondary fine hard martensite (αm*) and film-like γR) exists as the second phase. In addition, only a small amount of carbide (θ) precipitates only in the αm 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 fγ0 and Cγ0 are calculated by the methods proposed by Maruyama [82] and Dyson and Holmes [83], respectively.
On the other hand, the major phase of TM steel is αm 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 αm lath structure as the second phase [51,52,53]. The θ fraction (fθ) increases with decreasing TIT [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 αm 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 γR than TM steel [59].
The two-step Q&P process generally consists of quenching to a temperature (TQ) between MS and Mf after austenitizing and subsequent partitioning at a temperature (TP) higher than Ms (Figure 2b) [19,20,48]. The process forms the microstructure of αbf and αm matrix and γR, 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 TQ is illustrated in Figure 5a. On the quenching to TQ, a certain amount of austenite transforms to αm first. The αm fraction (fαm) can be estimated by the following empirical equation proposed by Koistinen and Marburger [84].
fαm = 1 − exp {−1.1 × 10−2 (MS − TQ)}
If the TQ is close to Mf, a small amount of carbide (θ) precipitates only in the αm lath structure [23,49]. During subsequent partitioning at temperatures above Ms, most of the remaining austenite transforms into αbf. At the same time, the αm 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 TQ dependence of fγ0 in two-step Q&P steels with different carbon content are shown in Figure 5b. In these two-step Q&P steels, the optimum TQ which gives the maximum volume fraction of γR is between Ms and Mf [19,20,23].
The mechanical properties of the third-generation AHSSs are predominantly controlled by the volume fraction and mechanical stability of γR. The mechanical stability is mainly related to its Cγ0, along with the size and stacking fault energy (SFE) of γR, matrix structure, deformation temperature, other microalloying elements, etc. [85]. According to Sugimoto et al. [86], the mechanical stability of γR can be defined by the following k-value (or the strain-induced transformation factor),
k = (ln fγ0 − ln fγ)/ε
where fγ is the retained austenite fraction after deformation to a certain plastic strain (ε). Sherif et al. [87] re-expresses the mechanical stability through the chemical free energy available for transformation as follows:
ln fγ0 − ln fγ = k1 ΔGα’γ ε
where k1 is a modified k-value. ΔGα’γ (= Gα’ − Gγ) 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 Gα’ and Gγ are the chemical free energies of ferrite (martensite) and austenite, respectively.
In general, the k- and k1-values decrease with increasing Cγ0 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 Cγ0, although the k-values are about three times higher than those of TPF and TAM steels because of lower Cγ0 [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 γR plays a role in the austenite stabilizer.

2.2. C-Si/Al-Mn Steel

In general, austenite-stabilizing elements such as Ni, Mn, C, N, and Cu, lower the critical temperature (T0) at which austenite and martensite have the same chemical free energy in steel [85]. Only Co increases the T0 and makes the austenite unstable. Many ferrite-stabilizing elements also stabilize the austenite, although they increase the T0. Exceptionally, Cr lowers the T0 and significantly increases the austenite stability by the addition of several mass% points, although it is a ferrite-stabilizing element [85].
Al makes the austenite unstable as a ferrite-stabilizing element, like the Co of an austenite-stabilizing element [85]. Ehrhardt et al. [90], Sugimoto and Mukherjee [69], and Alza and Chavez [91] summarize the effect of Al on the microstructure in TPF steel (Figure 7a). According to them, the advantage of Al over Si is increasing the driving force for austenite to bainite transformation, which accelerates the bainite transformation kinetics resulting from an increased nucleation rate. Examples of T0 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 T0 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 γR, 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 Ms [91].
For 0.2C-(0.5-1.5)Si-1.5Mn-(0.04-1.0)Al-(0-0.2)Mo-(0-0.05)Nb TBF steels [12] and 0.25C-(0.55-1.45)Si-(1.61-1.70)Mn-(0.3-0.69)Al Q&P steels [21], partial replacement of Si by Al increases the mechanical stability of γR and decreases its volume fraction. In this case, the increased mechanical stability is mainly associated with higher Cγ0 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 Cγ0 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.
The partial replacement of Si by Al refines the prior austenite grain, αbf lath, and film-like γR 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 αbf lath, film-like γR, 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 αbf 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 αm 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.
For the αm transformation, Kobayashi et al. [55] found that the αm 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 αbf 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 αm 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

According to Sugimoto et al. [15,54], flow stress (strain), σ(ε), of the AHSSs containing γR of 4 to 30 vol% is formulated by
σ(ε) = σM(ε) + Δσh(ε)
where σM(ε) and Δσh(ε) are the flow stress of the matrix and strain hardening increment of the steel, respectively. The Δσh(ε) can be estimated by
Δσh(ε) = Δσi(ε) + Δσt(ε) + Δσf(ε)
where Δσi(ε), Δσt(ε), and Δσf(ε) represent “the long-range internal stress hardening”, “the strain-induced transformation hardening”, and “the forest dislocation hardening”, respectively, which can be formulated by
Δσi(ε) = {(7−5ν)μ/5(1-ν)} f·εpu
Δσt(ε) = g(Δfαm)
Δσf(ε) = ζμ (b·f·ε/2r)1/2
where ν is the Poisson’s ratio, μ is the shear modulus, εpu is “the eigenstrain” [95], f is the volume fraction of the second phase, g(Δfαm) 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 γR. In the TBF, CFB, one-step Q&P, and BF-MMn steels subjected to the IT process at the temperatures above Ms, the second phase corresponds to untransformed carbon-enriched γR and strain-induced martensite. In these steels subjected to the IT process at the temperatures between Ms and Mf, the second phases correspond to αm, untransformed γR, 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].
Tensile ductility defined by uniform elongation (UEl), TEl, and reduction in area (RA) are mainly controlled by the strain-hardening behavior mentioned above in the third-generation AHSSs, which is related to the chemical composition and heat treatment conditions, such as austenitizing (annealing) temperature, TQ, TIT, TP, and these holding times. The effects of TQ 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 TQ = Ms − 100 °C, which was about 50 °C lower than TQ for the largest volume fraction of γR (Figure 4b). The optimum TQs 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 TQ for the TEl.
The effects of TIT 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 TIT above Ms, although the YS and TS considerably decrease. The optimum TIT for tensile ductility agrees well with one for the largest fraction of γR [51]. In the TM steel subjected to the IT process at temperatures below Mf, relatively high TS and TS×TEl are achieved, compared to the TBF steels subjected to the IT process at temperatures between Ms and Mf. 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.
For the TBF, CFB, one-step Q&P, two-step Q&P, BF-MMn, and Q&P-MMn steels, “the long-range internal stress hardening” and “the strain-induced transformation hardening” mainly contribute to the high ductility [10,14]. In the TM and M-MMn steels, a large amount of MA phase mainly increases “the long-range internal stress hardening”, with a small contribution to “the strain-induced transformation hardening” because of a small amount of γR or small Δfαm.
It is interesting that 0.2C-1.5Si-(1.5–5.0)Mn M-MMn [59] and D-MMn [39,40] steels achieve much higher TS×TEls than those of TBF and TM steels (Figure 10b).

3.2. C-Si/Al-MnSsteel

Partial replacement of Si by Al lowers the strain hardening rate and flow stress in 0.20C-(0.20-1.50)Si-1.24Mn-(0.022-1.22)Al-0.20Cr-(0.0026-0.0028)B TM steel (Figure 10a) [55,88], because solid solution-hardening of Al (24 MPa/at%) is about half that of Si (55 MPa/at%) in Fe-C ferrous steel [97]. Resultantly, the YS, TS, UEl, and TEl decrease with increasing Al content under a constant Si+Al content, although the RA increases with increasing Al content [88]. Additionally, a similar result can be obtained in 0.2C-(0.99-1.54)Si-1.51Mn-(0.033-0.49)Al-(0-0.05)Nb TBF steels subjected to the IT process at temperatures between Ms and Mf [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).
It is noteworthy that further addition of 0.05 mass% Nb is suitable to increase the TS×TEl in Al-bearing TBF steel [14] (see 0.5Al-0.05Nb TBF steel in Figure 10b). In this case, precipitation hardening by fine NbC contributes to increasing the TS. A similar result is reported in 0.2C-0.77Si-2.0Mn-0.76Al CFB steel [34], except for the TEl. In 0.25C-0.55Si-1.70Mn-0.69Al Q&P steel, 0.195C-4.52Mn-0.04Si-1.31Al Q&P-MMn steel, and 0.2C-0.08Si-4.04Mn-1.46Al Q&P-MMn steel. Low ductility resulting from the partial replacement of Si by Al was also reported by De Moor et al. [21,22], Kaar et al. [47], and Wallner et al. [49]. According to Sugimoto et al. [52] and Pham et al. [53], the volume fraction of θ in αm 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.
As shown in Figure 11a, the TS×TEls of various TBF and TM steels tend to increase with the volume fraction of γR [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 γR and low solid solution-hardening, although the mechanical stability of γR increases (or the k-value decreases). The TS×TEl− fγ0 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 γR, 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 γR.

4. Stretch formability

4.1. C-Si-Mn Steel

The stretch formability is defined by the maximum stretch height (Hmax), which is usually measured using a ball-head punch. Figure 12a shows the TIT dependence of the Hmax in 0.20C-1.50Si-1.50Mn-0.05Nb TBF and TM steels subjected to the IT process at TIT = 200 °C to 450 °C for 200 s after austenitizing [51]. For comparison, the TP dependence of the Hmax of 0.20C-1.50Si-1.50Mn-0.05Nb TM steel subjected to the DQ&P process after austenitizing and then partitioning at TP = 200 °C to 450 °C for 1000 s is shown in the Figure 12a. The Hmaxs values of the IT-processed TBF and TM steels (red line) increase with increasing TIT. The Hmax of DQ&P-processed TM steel increases with increasing TP, although the Hmaxs 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 γR. As shown in Figure 12b, the products of TS and Hmax (TS×Hmaxs) 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].
It is noteworthy that the Hmax 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 γR [39].

4.2. C-Si/Al-Mn Steel

In 0.20C-1.50Si-1.24Mn-0.20Cr-(0.022-1.22)Al-0.00bB TM steels (0Al, 0.7Al, and 1.2Al TM steels) under a condition of Si+Al = 1.5 mass%, the TS×Hmax 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 Hmax and TS×Hmax [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].
As shown in Figure 13b, the Hmax 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 Hmax and TS×Hmax of 1.2Al TM steel may be caused by lower UEl and Tel, resulting in low solution-hardening and low γR fraction. According to Sugimoto et al. [59], the small Hmax of the 5Mn M-MMn steel is caused by the presence of a much larger MA phase, although a large amount of metastable γR makes a positive contribution to the Hmax. 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).
The effect of a complex addition of Al and other alloying elements on the Hmax has not been investigated for the third-generation AHSSs. It is expected that the complex addition of Al and Nb enhances the Hmax because it achieves a large TEl in 0.5Al-0.05Nb TBF steel (Figure 10b).

5. Stretch-Flangeability

5.1. C-Si-Mn Steel

In general, the stretch-flangeability can be evaluated by the following hole-expansion ratio (HER) using a punched-hole specimen:
HER = (df − d0)/d0
where d0 and df 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.
The HER increases with increasing TIT in 0.20C-1.50Si-1.50Mn-0.05Nb IT-processed TBF and TM steels and increases with increasing TP in DQ&P-processed TM steel (Figure 14a) [51]. In this case, the TM steel subjected to the IT process at the temperatures of Mf (50 to 100) °C and TBF steel subjected to the IT process at the temperatures between Ms and Mf 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 αm lath structure with low θ fraction and (ii) plastic relaxation of localized stress concentration by the strain-induced transformation of metastable γR at the αm 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 αbf and αm, and the above (i) and (ii). The effects of TIT 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 TQ (in the Q&P process) on the stretch-flangeability was reported by Im et al. [25] using 0.18C-1.5Si-2.6Mn steel (Ms = 368 °C). In this case, higher HER was obtained by a lower quenching process at TQ (= 280 °C), followed by partitioning at TP = 425 °C.

5.2. C-Si/Al-Mn Steel

As shown in Figure 15a, partial replacement of Si by Al also keeps high TS×HER in 0.2C-(0.2-1.5)Si-1.24Mn-(0.022-1.22)Al-0.2Cr (0Al, 0.7Al, and 1.2Al) TM steels [55,88]. The TS×HER is at the same level as those of Al-free TM and TBF steels, such as 0.2C-1.5Si-1.5Mn-(0-1.0)Cr-(0-0.2)Mo (Cr-Mo) TM steels [53], 0.2C-1.0Si-1.5Mn-0.5Al (0.5Al) TBF steel [14], 0.2C-1.0Si-1.5Mn-0.5Al-0.05Nb (0.5Al-0.05Nb) TBF steel [14], and 0.2C-1.5Si-1.5Mn-0.05Nb (1.5Si) TBF steel [51]. De et al. [74], Sugimoto et al. [75], and Samek et al. [76] investigated the TS×HER of low-carbon Si/Al-Mn TPF [74,75,76] and TAM [75] steels. In the TPF and TAM steels, the replacement of Si by Al keeps a high TS×HER in the same way as the low-carbon Si/Al-Mn TBF and TM steels [14].
Differing from the drilled hole samples [69], the HER of the punching hole samples is considerably controlled by the surface-layer damage on hole-punching in TM and TBF steels [14,53,55]. The main surface-layer damage is measured by the size and number of voids and/or cracks in the break section on the punched surface, which are controlled by clearance between die [98] and punch, punching temperature and speed [89], and retained austenite characteristics [89]. In general, the punching damage can be also evaluated by a ratio of shear section length to sheet thickness (ss/t), because the larger the ss/t value, the higher the HER in TBF, TM, and M-MMn steels, as shown in Figure 16a. According to Kobayashi et al. [55], Sugimoto and Kobayashi [88], and Sugimoto et al. [14], the ss/t value increases with increasing Al content in 0Al, 0.7Al, and 1.2Al TM steels and 0.5Al-0.05Nb TBF steel. As shown in Figure 16b, the large ss/t of 1.2Al TM steel is connected to a large RA or local ductility. Therefore, the small punching damage may bring on high HER in 1.2Al TM steel, resulting from uniform αm lath structure, the high mechanical stability of γR, and low solution-hardening.
In 0.2C-(1.0-1.5)Si-1.5Mn-(0.04-0.5)Al-(0-0.05Nb) TBF steels subjected to the IT process at the temperatures just below Ms, 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 γR, and TRIP effect on hole expansion. In this case, the complex addition brings on precipitation-hardening by fine NbC/Mo2C.

6. Bendability

6.1. C-Si-Mn Steel

There is only a small amount of research on the bendability of the third-generation AHSSs [16,29,52,99]. The bendability is usually evaluated by the average bending angle [16,29] and minimum bending radius (Rmin) [52,99]. Figure 17a shows the variation in average bending angle as a function of austempering temperature (or TIT) 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 TIT = 350 °C, just higher than Ms (329 °C), as shown in Figure 17a. The optimum TIT corresponds to that of maximum total elongation (TE) (Figure 17b). In this case, the microstructure was a mixture of αbf and γR, and the γR 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 αm, and the high mechanical stability of γR.
Figure 18a shows the variations in Rmin 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 Rmin was obtained in the TBF steel subjected to the IT process at the temperatures between Ms and Mf and in the TM steel subjected to the IT process at 200 °C (<Mf = 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].
Figure 18b compares the Rmins of the TPF, TAM, and TBF steels with the chemistry of (0.1-0.6)C-1.5Si-1.5Mn [99]. Small Rmin values are achieved in TBF and TAM steels with lath-type uniform microstructure and high mechanical stability of γR, 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

Unfortunately, there are not any results on the effect of partial replacement of Si by Al on the bendability in the third-generation AHSSs. However, research investigating the effect of Al content on the Rmin in the first-generation AHSSs was reported by Sugimoto et al. [75]. Figure 19 shows the effect of Al content on Rmin − TIT 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 Rmin 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 γR needles. In general, small Rmin is brought about by large RA. As Al-added TBF and TM steels have large RA values, a small Rmin is expected to be achieved in the Al-added TBF and TM steels.

7. Summary and Perspectives

Partial replacement of Si by Al improves the coatability (or galvanizing property) of steel sheets, in the same way as P. Therefore, many researchers hope to know the effect of the partial replacement on the cold formability in low-carbon, Si-Mn, third-generation AHSSs, such as the TBF, Q&P, CFB, D-MMn, BF-MMn, and Q&P-MMn steels (Group I) and TM and M-MMn steels (Group II). The effects of the partial replacement of Si by 0.04 to 1.5 mass% Al on the microstructure, tensile properties, and cold formability of the AHSS sheets are summarized as follows.
(1) The partial replacement of Si by Al decreases the volume fraction of γR 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 γR fraction, resulting in decreased flow stress and a decreased strain hardening rate, despite increased mechanical stability of γR. Similar results are reported for CFB, Q&P, and Q&P-MMn steels.
(2) The partial replacement of Si by 1.2 mass% Al considerably decreases the stretch formability in the TM steel with TS above 1.2 GPa, although the stretch formability decreases with increasing Al content. This is associated with the fact that equi-biaxial tension growing on stretch-forming plays a role in deteriorating the stretch formability through easy crack and/or void formation. On the other hand, the partial replacement of Si by 0.5 to 1.2 mass% Al results in the same stretch-flangeability as Al-free steels in the TBF and TM steels. This was associated with small punching damage which results from a uniform lath structure and high mechanical stability of γR, despite low solution hardening and a decreased γR fraction. Unfortunately, there is not any research on the bendability in Al-added AHSSs.
(3) A complex addition of Al and Nb/Mo achieves considerably larger TEl and TS×TEl in TBF steels. In addition, the complex addition also further increases the stretch-flangeability of the TBF steel owing to the small punching damage due to refined microstructure, stabilized film-like γR, and TRIP effect on hole-expanding. In this case, the complex addition brings on precipitation hardening by fine NbC/Mo2C. 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.
(4) In order to apply the third-generation AHSSs to automotive sheet components, many studies on fatigue strength [6,54,81,101,102,103], toughness [6,29,43,54,57,81,88,104,105,106], hydrogen embrittlement resistance [13,18,107,108], and weldability [109,110,111,112,113,114,115,116,117,118,119] will be undertaken. It has been reported so far that the partial replacement of Si by Al increases the hydrogen embrittlement resistance in the 0.2C-(0.5-1.5)Si-1.5Mn-(0.04-1.0)Al TBF steel [13]. In the future, further research on the effects of the partial replacement of Si by Al on these mechanical properties of the AHSSs is expected. In parallel with this, many researchers try to apply the AHSSs to automotive hot/warm/cold forging parts [81,120,121,122,123,124]. In this field, further research investigating the effects of a complex addition of Al and other alloying elements on mechanical properties such as tensile properties, fatigue strength, toughness, and wear resistance is hoped for in the future.

Funding

This research received no external funding.

Institutional Review Board Statement

Not applicable.

Informed Consent Statement

Not applicable.

Conflicts of Interest

The author declares no conflict of interest.

Nomenclature

AHSSadvanced high-strength steelTRIPtransformation-induced plasticity
TWIPtwin-induced plasticityHMn TWIPhigh Mn TWIP
Aus.austeniticTBFTRIP-aided bainitic ferrite
Q&Pquenching and partitioningCFBcarbide-free bainite
D-MMnduplex type medium MnL-MMnlaminate type medium Mn
BF-MMnbainitic ferrite type medium MnQ&P-MMnQ&P type medium Mn
TMTRIP-aided martensiteM-MMnmartensite type medium Mn
TPFTRIP-aided polygonal ferriteTAMTRIP-aided annealed martensite
DPdual-phaseCPcomplex phase
Q&Tquenching and temperingDQ&Pdirect quenching and partitioning
ITisothermal transformationMsmartensite-start temperature
Mfmartensite-finish temperatureTITisothermal transformation temperature
TQquenching temperatureTPpartitioning temperature
T0critical temperature at which austenite and martensite have the same chemical free energy
γRretained austeniteαbfbainitic ferrite
αmprimary coarse soft martensiteαm*secondary fine hard martensite
MAMA (αm*+γR) phaseθcarbide
fγ0initial volume fraction of γRfγvolume fraction of γR
fαbfbainitic ferrite fractionfαmprimary martensite fraction
fαm*secondary martensite fractionfMAMA phase fraction
fαm’fαbf + fαm*fθcarbide fraction
Cγ0initial carbon concentration of γRεplastic strain
ΔGα’γchemical free energy change for transformation of γ to αGα’chemical free energy of ferrite (martensite)
Gγchemical free energy of austenitekstrain-induced transformation factor
k1modified k-valueSFEstacking fault energy
σflow stress of steelσMflow stress of matrix
Δσhstrain hardening incrementΔσilong-range internal stress
Δσttransformation hardeningΔσfforest dislocation hardening
νPoisson’s ratioμShear modulus
fvolume fraction of second phaseεpueigenstrain
Δfαmstrain-induced martensite fractionζmaterial constant
bBurgers vectorrparticle radius of second phase
YSyield stressTS, UTStensile strength
UEluniform elongationTEl, TEtotal elongation
RAreduction of areaTS×TElproduct of TS and TEl
Hmaxmaximum stretch heightTS×Hmaxproduct of TS and Hmax
HERhole expansion ratioTS×HERproduct of TS and HER
ss/tshear section length to sheet thicknessRminminimum bending radius

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Figure 2. Heat treatment diagrams of (a) TBF, CFB and TM steels [52] and (b) one-step and two-step Q&P steels [20]. TIT, TQ, TP, Ac3, Ms, and Mf are isothermal transformation temperature, quenching temperature, partitioning temperature, the austenite-finish temperature on heating, martensite-start temperature, and martensite-finish temperatures, respectively.
Figure 2. Heat treatment diagrams of (a) TBF, CFB and TM steels [52] and (b) one-step and two-step Q&P steels [20]. TIT, TQ, TP, Ac3, Ms, and Mf are isothermal transformation temperature, quenching temperature, partitioning temperature, the austenite-finish temperature on heating, martensite-start temperature, and martensite-finish temperatures, respectively.
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Figure 3. Illustration of the typical microstructure of various third-generation AHSSs [81]. (a): TBF, CFB, and BF-MMn steels (TIT > Ms); (b): TBF, CFB, one-step Q&P, and BF-MMn steels (TIT = Ms − Mf), and two-step Q&P and two-step Q&P-MMn steels (TP > Ms); (c): TM and M-MMn steels (TIT < Mf). αbf, αm, αm*, γR, θ, and MA represent bainitic ferrite, primary coarse soft martensite, secondary fine hard martensite, retained austenite, carbide, and MA phase (a mixture of αm* and film-like γR), respectively.
Figure 3. Illustration of the typical microstructure of various third-generation AHSSs [81]. (a): TBF, CFB, and BF-MMn steels (TIT > Ms); (b): TBF, CFB, one-step Q&P, and BF-MMn steels (TIT = Ms − Mf), and two-step Q&P and two-step Q&P-MMn steels (TP > Ms); (c): TM and M-MMn steels (TIT < Mf). αbf, αm, αm*, γR, θ, and MA represent bainitic ferrite, primary coarse soft martensite, secondary fine hard martensite, retained austenite, carbide, and MA phase (a mixture of αm* and film-like γR), respectively.
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Figure 4. Variations in initial volume fraction (fγ0, ●) and carbon concentration (Cγ0, ○) of retained austenite as a function of isothermal transformation temperature (TIT) 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 4. Variations in initial volume fraction (fγ0, ●) and carbon concentration (Cγ0, ○) of retained austenite as a function of isothermal transformation temperature (TIT) 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].
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Figure 5. (a) Illustration of variations in volume fractions of different phases as a function of quenching temperature (TQ) in two-step Q&P steel. (b) Variations in initial retained austenite fraction with TQ in 0.2C-4.0Mn-1.6Si-1.0Cr (Ms = 273 °C) and 0.3C-1.0Mn-1.6Si-1.0Cr (Ms = 235 °C) two-step Q&P steels [23]. (a) is reproduced based on references [19,20], in which fαm and fγ are volume fractions of primary coarse soft martensite and austenite as functions of TQ prior to partitioning. The final or initial austenite fraction (fγ0) at room temperature is given by a red, bold, solid line. fαm’ = fαbf (bainitic ferrite fraction) + fαm* (secondary fine hard martensite fraction). (b) is reprinted with permission from Elsevier, copyright 2022.
Figure 5. (a) Illustration of variations in volume fractions of different phases as a function of quenching temperature (TQ) in two-step Q&P steel. (b) Variations in initial retained austenite fraction with TQ in 0.2C-4.0Mn-1.6Si-1.0Cr (Ms = 273 °C) and 0.3C-1.0Mn-1.6Si-1.0Cr (Ms = 235 °C) two-step Q&P steels [23]. (a) is reproduced based on references [19,20], in which fαm and fγ are volume fractions of primary coarse soft martensite and austenite as functions of TQ prior to partitioning. The final or initial austenite fraction (fγ0) at room temperature is given by a red, bold, solid line. fαm’ = fαbf (bainitic ferrite fraction) + fαm* (secondary fine hard martensite fraction). (b) is reprinted with permission from Elsevier, copyright 2022.
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Figure 6. Relationships between strain-induced transformation factor (k) and initial carbon concentration of retained austenite (Cγ0) 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 6. Relationships between strain-induced transformation factor (k) and initial carbon concentration of retained austenite (Cγ0) 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.
Metals 12 02069 g006
Metals 12 02069 g007 550
Figure 7. (a) Effects of microalloying elements on the time-temperature transformation curve of steel [90]. (b) Calculated T0 curves and measured carbon concentration of retained austenite (Cγ0) 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 7. (a) Effects of microalloying elements on the time-temperature transformation curve of steel [90]. (b) Calculated T0 curves and measured carbon concentration of retained austenite (Cγ0) 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.
Metals 12 02069 g007
Metals 12 02069 g008 550
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 TP = 450 °C for 300 s after quenching [23]. Mss 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 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 TP = 450 °C for 300 s after quenching [23]. Mss 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.
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Metals 12 02069 g009 550
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 (TIT) in 0.20C-1.50Si-1.50Mn-0.05Nb TBF and TM steels. This figure is reproduced based on reference [51].
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 (TIT) in 0.20C-1.50Si-1.50Mn-0.05Nb TBF and TM steels. This figure is reproduced based on reference [51].
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Metals 12 02069 g010 550
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 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].
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Metals 12 02069 g011 550
Figure 11. (a) Relationship between tensile strength (TS) and initial retained austenite fraction (fγ0). (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 11. (a) Relationship between tensile strength (TS) and initial retained austenite fraction (fγ0). (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].
Metals 12 02069 g011
Metals 12 02069 g012 550
Figure 12. (a) Isothermal transformation (TIT) and partitioning temperature (TP) dependences in maximum stretch height (Hmax) 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 Hmax 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 12. (a) Isothermal transformation (TIT) and partitioning temperature (TP) dependences in maximum stretch height (Hmax) 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 Hmax 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].
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Metals 12 02069 g013 550
Figure 13. (a) Maximum stretch height–tensile strength (Hmax−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 Hmax 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 13. (a) Maximum stretch height–tensile strength (Hmax−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 Hmax 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].
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Metals 12 02069 g014 550
Figure 14. (a) Isothermal transformation and partitioning temperature (TIT and TP) 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 14. (a) Isothermal transformation and partitioning temperature (TIT and TP) 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].
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Metals 12 02069 g015 550
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 (TIT) 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 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 (TIT) 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.
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Metals 12 02069 g016 550
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 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].
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Metals 12 02069 g017 550
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]. Ms of the steel is 329 °C.
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]. Ms of the steel is 329 °C.
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Metals 12 02069 g018 550
Figure 18. Relationships between minimum bending radius (Rmin) 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 18. Relationships between minimum bending radius (Rmin) 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.
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Metals 12 02069 g019 550
Figure 19. Variations in minimum bending radius (Rmin) as a function of isothermal transformation temperature (TIT) 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.
Figure 19. Variations in minimum bending radius (Rmin) as a function of isothermal transformation temperature (TIT) 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.
Metals 12 02069 g019
Table
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.
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.SteelChemical CompositionPropertyRef.
1st
Gen.		TPF0.25C-1.28Si-1.67Mn-0.03Al, 0.18C-0.02Si-1.56Mn-1.73Al1[70]
0.21C-2.10Si-1.52Mn-0.022Al, 0.22C-0.01Si-1.49Mn-2.02Al1, 2[71]
0.21C-2.10Si-1.52Mn-0.022Al, 0.22C-0.01Si-1.49Mn-2.02Al1, 2, 4[72]
0.19C-1.46Si-1.57Mn-0.06Al, 0.31C-0.34Si-1.57Mn-1.23Al1, 2[73]
(0.14-0.21)C-(0.34-1.47)Si-1.5Mn-(0.03-0.99)Al1, 2[74]
0.20C-(0.49-1.54)Si-1.5Mn-(0.04-0.99)Al1, 2, 4, 5[75]
(0.19-0.25)C-(0.09-1.45)Si-1.7Mn-(0.03-1.49)Al1[76]
0.20C-1.87Si-1.99Mn-(0.04-2.0)Al1[77]
0.20C-(0.49-1.50)Si-1.5Mn-(0.04-0.99)Al1, 2, 4[78]
TAM0.20C-(0.49-1.54)Si-1.5Mn-(0.04-0.99)Al1, 2, 4, 5[75]
0.20C-(0.48-1.50)Si-1.5Mn-(0.04-0.99)Al1, 2, 4[78]
3rd Gen.TBF0.20C-(0.49-1.54)Si-(1.48-1.51)Mn-(0.04-0.99)Al-(0-0.05)Nb-(0-0.20)Mo1, 2, 4[12]
0.20C-(0.49-1.51)Si-(1.51-2.51)Mn-(0.04-0.99)Al1, 2[13]
0.20C-(0.99-1.54)Si-1.5Mn-(0.04-0.49)Al-(0-0.05)Nb1, 2, 4[14]
Q&P0.24C-1.45Si-1.61Mn-0.30Al, 0.25C-0.55Si-1.70Mn-0.69Al1,2[21]
0.24C-0.12Si-1.60Mn-1.41Al-0.17Mo2[22]
0.30C-(0.48-0.99)Si-(1.86-2.00)Mn-(0.01-1.10)Al-(1.01-2.20)Cr1[24]
CFB0.25C-(0.08-1.09)Si-2.07Mn-(0.021-1.54)Al1, 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-0.23Mo1, 2[32]
0.2C-1.55Si-2.0Mn, 0.2C-0.77Si-2.0Mn-0.76Al1, 2[34]
D-MMn(0.1-0.3)C-(0-1.5)Si-(2-5)Mn-(0-1.5)Al-(0-1.5)Cr1[38]
(0.18-0.19)C-(7.66-7.93)Mn-(0-2.79)Al1, 2[42]
BF-MMn0.18C-0.23Si-3.6Mn-1.7Al-0.2Mo-0.04Nb1[45]
Q&P-MMn0.173C-4.46Mn-1.47Si-0.03Al, 0.195C-4.52Mn-0.04Si-1.31Al1, 2[47]
0.2C-1.50Si-4.02Mn-0.02Al, 0.2C-0.08Si-4.04Mn-1.46Al1, 2[49]
TM0.20C-(0.20-1.50)Si-1.24Mn-(0.02-1.22)Al-0.2Cr-(0.003-0.005) Ti-(0.003-0.005)B1, 2, 3, 4[55]
1: microstructure, 2: tensile properties, 3: stretch formability, 4: stretch-flangeability, 5: bendability.
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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

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