ISIJ International
Online ISSN : 1347-5460
Print ISSN : 0915-1559
ISSN-L : 0915-1559
Regular Article
Fracture Toughness of an Advanced Ultrahigh-strength TRIP-aided Steel
Junya KobayashiDaiki InaAsahiko FutamuraKoh-ichi Sugimoto
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2014 Volume 54 Issue 4 Pages 955-962

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Abstract

The fracture toughness of an advanced ultrahigh-strength 0.2%C-1.5%Si-1.5%Mn-1.0%Cr-0.05%Nb (in mass%) transformation-induced plasticity (TRIP)-aided steel with a bainitic ferrite and/or martensite structure matrix was investigated for applications in automobiles, construction machines, and pressure vessels. After the steel was austenitized and isothermally transformed via heat treatment at temperatures between 200°C and 350°C below the martensite-finish temperature, it exhibited a good combination of tensile strength (1.4 GPa) and total elongation (15%). In addition, the steel achieved a much higher plane-strain fracture toughness (KIC = 129–154 MPa m1/2) than conventional structural steel such as SCM420 steel (KIC = 57–63 MPa m1/2). Surprisingly, the fracture toughness was nearly the same as that of a maraging steel. Our results indicate that the high fracture toughness was associated with (1) a softened wide lath-martensite matrix with a low carbide content and carbon concentration and (2) effective plastic relaxation of localized stress concentration by the strain-induced transformation of fine metastable retained austenite in the narrow lath-martensite and retained austenite mixture, which suppresses void formation and cleavage crack initiation at the pre-crack tip.

1. Introduction

Ultrahigh-strength transformation-induced plasticity (TRIP)-aided steels with bainitic ferrite and/or martensite structure matrix1,2,3) have recently been developed as third-generation ultrahigh-strength cold-rolled sheets for automotive bodies. In addition, other ultrahigh-strength steels such as quench and partitioning (Q&P) steel4,5) and 5–25%Mn TRIP/twining-induced plasticity (TWIP) steels6,7) were also developed. These steels possess excellent elongation, stretch-flangeability and bendability owing to the TRIP effect of metastable-retained austenite, and the steels have thus been partially used for the structural parts of automobiles. TRIP-aided steels also promise to be next-generation structural steels in the forging of mechanical parts of various machines because of their excellent impact toughness8,9) and high notch fatigue strength10,11) resulting from the effect of plastic relaxation on the strain-induced transformation of metastable-retained austenite of 3–10 vol%.

If these TRIP-aided steels are to be used in higher pressure vessels and further high strength construction machines, high fracture toughness is also required. However, there has been no report on the fracture toughness of ultrahigh-strength TRIP-aided steels.

The current study investigated the fracture toughness of ultrahigh-strength TRIP-aided steels produced by isothermal transformation (IT) at different temperatures in linear and nonlinear fracture toughness tests employing acoustic emission (AE). In addition, the fracture toughness was related to metallurgical properties such as the microstructural and retained-austenite characteristics, and the strain-induced transformation behavior of the retained austenite.

2. Experimental Procedure

Vacuum-melted and hot-forged slabs were prepared with the chemical composition given in Table 1. These slabs were then hot-rolled to produce plates of 15 mm thickness and bars of 13 mm diameter. The martensite-start (Ms) and -finish (Mf) temperatures were respectively 401°C and 261°C, where Ms and Mf were obtained from a dilatation curve at a cooling rate of 30°C/s. Ms was determined by the temperature at which the specimen was expanded by martensite transformation. Mf was determined by the temperature at which the expansion of specimen stopped. Nb was added to refine the prior austenitic grain. Cr was also added to improve the hardenability of the steel. For comparison, commercial SCM420 steel bars of 36 mm diameter were hot-rolled to make plates of 15 mm thickness and bars of 13 mm diameter.

Table 1. Chemical composition (mass%) and measured martensite-start (Ms) and -finish temperatures (Mf, °C) of steels used.
SteelCSiMnPSCuNiCrMoAlNbONMsMf
TRIP0.211.491.500.00190.004< 0.02< 0.021.00< 0.010.0400.0500.00140.0013401261
SCM4200.210.210.770.02000.0240.110.061.020.18******
*:  not measured

Tensile specimens (JIS 14A: 5 mm diameter by 20 mm gauge length) and compact specimens (ASTM-E399: 24 mm width and 12 mm thickness) were machined from these bars and plates, respectively (Fig. 1). The specimens were subjected to IT process at isothermal transformation temperature (TIT) = 50°C–450°C for isothermal transformation time (tIT) = 1000 s after austenitizing at 900°C, as shown in Fig. 2.

Fig. 1.

(a) Dimensions and (b) machining directions of the compact specimen.

Fig. 2.

Heat treatment diagram of TRIP-aided steel, in which αbf, αm, αm*, MA, and γR denote bainitic ferrite, wide lath-martensite, narrow lath-martensite, MA-like phase, and retained austenite respectively.

The microstructure of the steels was observed by scanning electron microscopy (SEM), transmission electron microscopy (TEM) and field-emission-SEM/electron backscatter diffraction (EBSD) analysis. The volume fraction of retained austenite (fγ, vol%) was calculated from the integrated intensity of the (200)α, (211)α, (200)γ, (220)γ and (311)γ peaks obtained by X-ray diffractometry using Mo-Kα radiation.12) The carbon concentration (Cγ, mass%) was estimated using Eq. (1). In this case, the lattice constant (aγ, units of 10−1 nm) of retained austenite was calculated from the (200)γ, (220)γ and (311)γ peaks of Cu-Kα radiation.13)   

a γ =3.5780+0.0330% C γ +0.00095%M n γ -0.0002%N i γ +0.0006%C r γ +0.0220% N γ +0.0056%A l γ -0.0004%C o γ +0.0015%C u γ +0.0031%M o γ +0.0051%N b γ +0.0039%T i γ +0.0018% V γ +0.0018% W γ , (1)
where %Mnγ, %Niγ, %Crγ, %Nγ, %Alγ, %Coγ, %Cuγ, %Moγ, %Nbγ, %Tiγ, %Vγ and %Wγ denote the concentrations of the respective individual elements (mass%) in the retained austenite. As an approximation, the contents of added alloying elements were substituted for these concentrations in this study. The volume fraction of carbide in the specimens was measured as follows. First the specimens were etched in 5%-nital etchant, and then carbon was coated on the specimen. Secondly, the carbon reprica with carbides was peeled off from the specimen in ethanol solution including 30% nitric acid, followed by TEM examination and image analysis in an area of about 560 μm2.

Tensile tests were carried out on a tensile testing machine (Shimadzu Co., AG-10TD) at 25°C and a crosshead speed of 1 mm/min (a strain rate of 8.33 × 10−4 s–1). Fracture toughness tests were conducted on the same testing machine at 25°C using a clip gauge, after a pre-crack was introduced in a fatigue test according to ASTM and JSME standards.14,15) Simultaneously, the AE technique was applied to detect a crack initiation signal and strain-induced transformation behavior of retained austenite using a PCI-2 system (PAC, USA).16) The testing system is shown in Fig. 3. For this purpose, a wideband differential piezo-electric transducer (WD, 18 mm diameter) with an operating frequency range of 100–1000 kHz was used. A wideband sensor was selected in preference to a resonant sensor because resonant sensors detect only a narrow band of frequencies. To ensure good coupling, lead break tests were performed before all tests using pencil leads with diameter of 0.5 mm. The transducer was connected to a preamplifier with gain offset to 40 dB. To remove noise from the actuator and other components, a threshold amplitude of 30 dB was selected after performing preliminary tests.

Fig. 3.

Setup of the fracture toughness test.

3. Results

3.1. Microstructure and Tensile Properties

Figure 4 shows SEM images of TRIP-aided steels isothermally transformed at temperatures between 50°C and 450°C and SCM420 steel tempered at 200°C. EBSD analysis results of TRIP-aided steel are shown in Fig. 5. In Figs. 5(a) and 5(d), orange, yellowish green and black regions show the matrix structure, martensite–austenite complex phase and retained austenite, respectively. The volume fraction of blocky second phase was measured from the yellowish green region on image quality distribution map by EBSD analysis. In this study, it was assumed that the region of IQ index < 4675 is blocky second phase. This yellowish green region was confirmed to be almost the same as the blocky second phase observed by SEM. Hereafter, the above complex phase is called the MA-like phase, because it resembles the MA constituent in conventional bainitic steel. When IT process was conducted at temperatures (i) higher than the Ms temperature (401°C), the microstructure of the TRIP-aided steel principally changed to bainitic ferrite, MA-like phase and isolated retained austenite as shown in Figs. 4(f), 4(g) and 5(d). When IT process was carried out at temperatures (ii) between the Ms (401°C) and Mf (261°C) temperatures, the microstructure mainly consisted of bainitic ferrite, wide lath-martensite, finely dispersed MA-like phase and isolated retained austenite (Figs. 4(d) and 4(e)). In contrast, when IT process was carried out at temperatures (iii) below the Mf temperature (261°C), the microstructure changed to wide lath-martensite and finely dispersed MA-like phase on prior austenitic grain, packet and martensite block boundaries (Figs. 4(a)–4(c) and 5(a)). The above results agree well with the previous studies2,3,9) Thus hereafter, these TRIP-aided steels with microstructures obtained in cases (i), (ii) and (iii) above are respectively called TRIP-aided bainitic ferrite (TBF), TRIP-aided bainitic ferrite/martensitic (TBM) and TRIP-aided martensitic (TM) steels. SCM420 steel consisted of a wide lath-martensite structure and MA-like phase (Fig. 4(h)). However, the MA-like phase contained little retained austenite.

Fig. 4.

SEM images of TRIP-aided steel subjected to IT process at (a) TIT = 50, (b) 200, (c) 250, (d) 300, (e) 350, (f) 400 and (g) 450°C and (h) SCM420 steel tempered at TT = 200°C, in which αbf, αm and MA denote bainitic ferrite, martensite and MA-like phase, respectively.

Fig. 5.

Image quality distribution maps of body centered cubic lattice (bcc) and orientation maps of bcc and face centered cubic lattice (fcc) in TRIP-aided steel subjected to IT process at TIT = 200 and 400°C. αm, αm*, αbf, MA and γR denote wide lath-martensite, narrow lath-martensite, bainitic ferrite, MA-like phase and retained austenite, respectively.

In the TBF and TBM steels, the retained austenite phases are mainly in the blocky MA-like phase or along the bainitic ferrite lath boundary (Fig. 5(d)). The volume fraction of martensite (fαm) increases with decreasing quenching temperature (in this study, quenching temperature is isothermal transformation temperature (TIT)) according to an equation reported by Koistinen and Marburger:17)   

f α m =1-exp{-1.1× 10 -2 ( M S - T IT )}, (2)
In this study, the volume fraction of wide lath-martensite in the TRIP-aided steel also increased with decreasing isothermal transformation temperature (Figs. 4 and 5). For the TRIP-aided steel, most of the MA-like phase is located on prior austenitic grain, packet and block boundaries. The MA-like phase fraction tends to be slightly higher in the TRIP-aided steel subjected to IT process at lower temperature, except for TIT = 450°C (Table 2). SEM-EBSD analysis showed that the SCM420 steel tempered at 200°C contained the same MA-like phase fraction as TRIP-aided steel, although most of the MA-like phase was formed only by the narrow lath-martensite structure.
Table 2. Retained austenite characteristics, tensile properties and fracture toughness of TRIP-aided and SCM420 steels.

TIT (°C): isothermal transformation temperature, TT (°C): tempering temperature, fγ0 (vol%): initial volume fraction of retained austenite, Cγ0 (mass%): initial carbon concentration of retained austenite, k: strain-induced transformation factor, fMA: volume fraction of the MA-like phase, fθ: volume fraction of carbide, YS (MPa): yield stress or 0.2% offset proof stress, TS (MPa): tensile strength, YR: yield ratio (= YS/TS), UEl (%): uniform elongation, TEl (%): total elongation, RA (%): reduction of area, Kin (MPa m1/2): fracture toughness upon initial cracking, KQ (MPa m1/2): provisional fracture toughness, Jin (MPa m): J-integral value upon crack initiation, JIC (MPa m): J-integral value near the onset of stable crack extension.

Figure 6 shows TEM images of the TRIP-aided steels isothermally transformed at 200°C and 400°C. When the steel was subjected to IT process at temperatures lower than 350°C, a small amount of fine carbides precipitated in the wider lath-martensite structure (Fig. 6(a)). In contrast, there was no carbide in the steel subjected to IT process at 400°C and 450°C. The amount of carbide increased with decreasing TIT, as shown in Table 2.

Fig. 6.

TEM images of TRIP-aided steel subjected to IT process at (a) TIT = 200 and (b) 400°C, in which αm, αbf, γR and θ denote wide lath-martensite, bainitic ferrite, retained austenite and carbide, respectively.

Figure 7 shows the retained-austenite characteristics of TRIP-aided steels isothermally transformed at temperatures between 50°C and 450°C. The volume fraction of retained austenite in SCM420 steel is listed in Table 2. The volume fraction of retained austenite increases with increasing TIT, except for a sharp decrease at 450°C. The variation in the carbon concentration is complex, with the minimum and maximum concentrations observed at TIT = 200°C and 350°C, respectively.

Fig. 7.

Variations in (a) initial volume fraction (fγ0) and carbon concentration (Cγ0) of retained austenite and (b) k value as a function of isothermal transformation temperature (TIT) in TRIP-aided steel.

Figure 7(b) shows the strain-induced transformation factor k for TRIP-aided steel, defined by   

log f γ = log f γ0 -kε, (3)
where fγ0 and fγ are the volume fractions of retained austenite before and after straining to the plastic strain ε, respectively. In general, a lower k value indicates higher mechanical stability of retained austenite and is related to a higher concentration of carbon in the retained austenite.1,2) SCM420 steel tempered at 200°C contained only a little retained austenite (1.10 vol%) in the MA-like phase, and the amount of retained austenite considerably decreased with increasing tempering temperature (Table 2).

The tensile properties of the TRIP-aided steels are given in Table 2 and Fig. 8. If the steel is subjected to IT process at temperatures between 50°C and 350°C, higher yield stress or 0.2% offset proof stress (YS = 1121–1177 MPa) and tensile strength (TS = 1405–1599 MPa) are achieved. The tensile strength increases with decreasing TIT, although the yield stress is hardly affected by the TIT. When compared with SCM420 steel, the TRIP-aided steels have lower yield ratios and larger total elongations and greater reduction of areas (Table 2).

Fig. 8.

Variations in (a) yield stress or 0.2% offset proof stress (YS) and tensile strength (TS) and (b) uniform (UEl), local (LEl) and total elongations (TEl) as a function of the isothermal transformation temperature (TIT) in TRIP-aided steel.

3.2. Fracture Toughness

Figure 9 shows load–displacement curves of TRIP-aided and SCM420 steels with AE signals. Both steels exhibited ductile fracture behavior. The maximum load and displacement to fracture of TRIP-aided steel isothermally transformed at 200°C are larger than those of TRIP-aided steel isothermally transformed at 400°C or SCM420 steel. AE signals in TRIP-aided steel originated before initial cracking, differing from the case for SCM420 steel, which means that some of the retained austenite was transformed by strain before initial cracking, which plastically relaxed the localized stress concentration around the crack tip.

Fig. 9.

Load–displacement (Pδ) curves and AE signals of TRIP-aided steel subjected to IT process at (a) TIT = 200 and (b) 400°C and (c) SCM420 steel tempered at TT = 200°C.

Figure 10 shows the variation in provisional fracture toughness (KQ) as a function of TIT for the TRIP-aided steels. The provisional fracture toughness was calculated as14)   

K Q = P Q f( α ) /B W 1/2 , (4)
where f(α) = 29.6α1/2 − 185.5α3/2 + 655.7α5/2 − 1017.0α7/2 + 638.8α9/2 and α = a/W. Here PQ is the provisional applied load and a, B and W are respectively the crack length, thickness and width of a compact specimen as seen in Fig. 1. KQ values of the TRIP-aided steel isothermally transformed at TIT = 200 through 350°C are higher than values for steels subjected to IT process at other temperatures and nearly equal to the fracture toughness upon stable crack initiation (Kin) calculated from AE (Table 2). If both the following Eqs. (5) and (6) are satisfied, the calculated provisional fracture toughness KQ can be evaluated as the plane-strain fracture toughness KIC.14)   
P max / P Q < 1.10, (5)
  
a,B> 2.5 K Q 2 /Y S 2 , (6)
where Pmax is the maximum load (see Fig. 9(a)). In this study, the specimen thickness (B = 12 mm) of TRIP-aided steels subjected to IT process at 50°C–350°C was less than B in Eq. (6), as for SCM420 steel tempered at 400°C (Table 2).
Fig. 10.

Variation in fracture toughness (KQ) as a function of the isothermal transformation temperature (TIT) in TRIP-aided steel.

Figure 11 shows the blunting lines and R-curves of typical TRIP-aided and SCM420 steels. Since the J-integral value upon crack initiation (Jin), which is determined from the intersection of the blunting line and R-curve, is satisfied by Eq. (7), Jin can be considered equal to JIC (the value of the J-integral near the onset of stable crack extension).15)   

B,b>25 J in / σ fs , (7)
where b is the length of the uncracked ligament of the specimen (Wa) and σfs = (YS + TS) / 2.
Fig. 11.

Blunting line and R-curves of TRIP-aided steel subjected to IT process at TIT = 200°C and SCM420 steel tempered at TT = 200°C, in which σfs = (YS + TS)/2 and Δa is the crack extension measured from the exposed fracture surface.

When the plane-strain fracture toughness KIC can be calculated by Eq. (8),15) the estimated values of KIC (137 MPa m1/2 for TRIP-aided steel, 67 MPa m1/2 for SCM420 steel) are nearly the same as the experimental values of KQ (Table 2). Therefore, the KQ values in Fig. 10 can be regarded as KIC values.   

J IC = (1-ν) K IC 2 /E, (8)
where ν is Poisson’s ratio (0.28) and E is Young’s modulus (206 GPa).

Figure 12 shows the relationship between provisional fracture toughness (KQ) and yield stress in TRIP-aided and SCM420 steels. The figure includes KQ values of TRIP-aided steels subjected to IT process at 50°C and subsequent partitioning process (heat treatment for carbon enrichment of retained austenite) at 50°C through 350°C (ITP process),18) and KIC values of 18Ni maraging steel19,20) and other steels.21) Note that the present TRIP-aided steels subjected to IT process are characterized by the same level of fracture toughness as the maraging steel, although they exhibit much lower fracture toughness than the high-alloy TRIP steel.

Fig. 12.

Relationship between fracture toughness (KQ or KIC) and yield stress (YS) in TRIP-aided (●: IT process, ○: ITP process),18) SCM420 (□), 18Ni maraging,19,20) Fe–C–Cr–Ni–Mn TRIP, AISI4340 and other steels.21)

Figure 13 shows SEM images of the fracture surfaces in the TRIP-aided and SCM420 steels. A distinct ductile fracture is seen only in the TRIP-aided steel subjected to IT process at 200°C–350°C (Fig. 13(a)) and the SCM420 steel tempered at 200°C (Fig. 13(c)). It is noteworthy that the former ductile fracture consists of fine and coarse dimples. In addition, the spacing between coarse dimples is 10–20 μm and equivalent to that between larger MA-like phases. On the other hand, there was cleavage fracturing on the fracture surface of TRIP-aided steels subjected to IT process at 400°C and 450°C (Fig. 13(b)) and SCM420 steels tempered at 300°C and 400°C (Fig. 13(d)).

Fig. 13.

SEM images of the fracture surface of TRIP-aided steel subjected to IT process at (a) TIT = 200 and (b) 400°C and SCM420 steel tempered at (c) TT = 200 and (d) 300°C, in which SZW represents stretched zone width.

4. Discussion

4.1. Optimum Microstructure to Obtain High KQ

In general, the fracture toughness of TRIP-aided steel is expected to be controlled in the same way as the impact toughness— by the matrix structure and characteristics of retained austenite (volume fraction, carbon concentration and morphology), the MA-like phase (morphology, volume fraction and site) and carbide fraction.9) Figure 14 shows the relationships between KQ and retained austenite characteristics in SCM420 steel tempered at 200°C, 300°C and 400°C and TRIP-aided steel subjected to IT process at temperatures between 50°C and 350°C, where the carbon concentration of retained austenite in SCM420 steel was assumed to be 0.2 mass%. From these results, we conclude that high KQ values can be obtained with a high volume fraction, high carbon concentration and high stability of the retained austenite. The relationship between fracture toughness and k in Fig. 14(c) agrees well with a report by Gerberich et al.22) that there is the following relationship between plane stress fracture toughness KC and transformation coefficient m for C–Cr–Ni–Mo–Mn–Si system TRIP steels,   

K C m 1/2 (9)
where m is related to the alloy stability in terms of the strain-induced phase transformation, and is similar to k in this study. It is known that high fracture toughness of maranging steel results from low carbon and high alloying matrix strengthened by fine precipitates.19,20,23) If the maraging steel was subjected to over aging, a large amount of reverted austenite can be formed. In this case, the reverted austenite plays a role in further increasing fracture toughness, similar to the present TBM and TM steels.
Fig. 14.

Relationships between KQ and (a) volume fraction (fγ0), (b) carbon concentration (Cγ0) and (c) strain-induced transformation factor (k) of retained austenite in TRIP-aided (●) and SCM420 steels (□).

When the TRIP-aided steel was isothermally produced at 400°C or 450°C, KQ decreased considerably. In Figs. 4(f), 5(d) and 6(b), the matrix structure of the TRIP-aided steel isothermally transformed at 400°C was a coarse bainitic ferrite lath-structure. In addition, a large amount of blocky MA-like phase had formed, although a single phase of retained austenite was also present. Since the blocky MA-like phase is expected to behave like a stress concentration site, the low KQ value of TRIP-aided steels may be due to a coarse matrix structure and blocky MA-like phase and a large difference in flow stress between the matrix and MA-like phase. On the other hand, higher KQ values in the cases of 200°C–350°C are considered to be due to a fine mixture of wide lath-martensite and metastable retained austenite of about 5 vol% in the MA-like phase. In this case, the strength ratio of MA-like phase to matrix is relatively low, and void formation at the interface between MA-like phase and matrix is thus considerably suppressed, as opposed to the case at 400°C or 450°C.

4.2. Improvement Mechanism of KQ

In Fig. 13(a), a ductile fracture consisting of fine and coarse dimples appeared on the fracture surface of the TRIP-aided steel subjected to IT process at 200°C, with a wide stretched zone width. Because the fracture is assumed to follow the Rice and Johnson model,24) namely crack blunting, void initiation and growth, and void coalescence, the fracture behavior can be illustrated as shown in Fig. 15. In general, the initiation, growth and coalescence behavior of voids in bainitic or martensitic steels is controlled by interparticle paths of the second phases (MA-like phases) and carbides. According to Horn and Ritchie25) and Sarikaya et al.,26) the carbide precipitates are located on the prior austenitic grain, packet, and block boundaries, as well as in the wide lath-martensite structure. These carbide precipitates generally act as void-initiation sites in martensitic steels. In the present TBM and TM steels, however, a small number of voids primarily initiate only at the larger MA-like phase/wide lath-martensite structure matrix interface, not at matrix/carbide interface. At the larger MA-like phase/wide lath-martensite structure matrix interface at which there is a highly localized stress concentration results in, the voids then grow into the coarse dimples with the coalescence of the coarse voids and the forming of fine dimples (Fig. 15). In addition, the void formation is affected by the strain-induced transformation of the retained austenite because some retained austenite phases are always included in the plastic region (dY) of the pre-crack tip estimated as27)   

d Y = K 2 /(3πY S 2 ), (10)
where K is a stress intensity factor defined as σ(πc)1/2, σ is the applied stress, and c is the crack length. For steel isothermally produced at 200°C–350°C, if K and YS in Eq. (10) are Kin and YS listed in Table 2, dY is calculated as 1.2–1.7 mm.
Fig. 15.

Illustration of a ductile fracture consisting fine and coarse dimples in TRIP-aided steel subjected to IT process at TIT = 200°C, in which αm, αm*, MA, γR and θ denote wide lath-martensite, narrow lath-martensite, MA-like phase, retained austenite and carbide, respectively.

According to the work of Kobayashi et al.9) on tensile deformation of TM steel, most voids are formed at the MA-like phase/matrix interface and the strain-induced transformation of the retained austenite in MA-like phase makes it difficult for voids to form because it relaxes plastically the localized stress concentration at the interface. Furthermore, it makes void coalescence or extension difficult because of the wide space between the large voids. The current TRIP-aided steel isothermally produced at 200°C–350°C contained highly stable retained austenite and a large amount of finely dispersed MA-like phase, as well as a softened wide lath-martensite structure matrix containing little carbide. Therefore, it is considered that the void formation is primarily disturbed by (i) the plastic relaxation resulting from the strain-induced transformation of metastable-retained austenite and (ii) a softened wide lath-martensite structure matrix, as well as a small number of void initiation sites. These lead to higher fracture toughness through the difficulty of voids to coalesce compared with the case for SCM420 steel.

5. Summary

The fracture toughness of 0.2%C-1.5%Si-1.5%Mn-1.0%Cr-0.05%Nb ultrahigh-strength TRIP-aided steel with bainitic ferrite and/or martensite structure matrices was investigated to develop the next generation of structural steel. Important results are summarized as follows.

(1) When the TRIP-aided steel was isothermally transformed at temperatures between 200°C and 350°C below Ms after austenitizing, the steel had a tensile strength of 1.4 GPa and total elongation of 15%.

(2) The steel achieved a much higher fracture toughness (129–154 MPa m1/2) than conventional structural SCM420 steel quenched and tempered at 200°C or 300°C (57–63 MPa m1/2). The fracture toughness was the same as that of maraging steel.

(3) It is considered that the superior fracture toughness is essentially due to (i) a softened wide lath-martensite matrix containing little carbide and (ii) effective plastic relaxation by the strain-induced transformation of fine metastable-retained austenite in the finely dispersed MA-like phase, which suppresses void formation, growth and coalescence as well as cleavage fracture at the pre-crack tip.

Acknowledgment

This study was supported by grants from the Adaptable and Seamless Technology Transfer Program through Target-driven R&D from the Japan Science and Technology Agency.

References
 
© 2014 by The Iron and Steel Institute of Japan
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