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Nanoscale cementite and microalloyed carbide strengthened Ti bearing low carbon steel plates in the context of newly developed ultrafast cooling
Materials Science & Engineering A 698 (2017) 268–276
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Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Nanoscale cementite and microalloyed carbide strengthened Ti bearing low carbon steel plates in the context of newly developed ultrafast cooling
MARK
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X.-L. Lia, C.-S. Leia, Q. Tianb, X.-T. Denga, , L. Chenb, P.-L. Gaob, K.-P. Duc, Y. Dud, Y.-G. Yuc, ⁎⁎ Z.-D. Wanga, , R.D.K. Misrae a
State Key Laboratory of Rolling and Automation, Northeastern, Shenyang 110819, China Key Laboratory of Neutron Physics and Institute of Nuclear Physics and Chemistry, China Academy of Engineering Physics, Mianyang 621999, China c Beijing General Research Institute of Mining and Metallurgy, Beijing 100160, China d Materials Science and Engineering, Northwestern University, Evanston 60208, USA e Laboratory for Excellence in Advanced Steel Research, Department of Metallurgical, Materials and Biomedical Engineering, University of Texas at El Paso, El Paso, TX 79968-0521, USA b
A R T I C L E I N F O
A B S T R A C T
Keywords: High strength low alloy Thermomechanical controlled processing Small-angle neutron and X-ray scattering Microstructural evolution High resolution transmission electron microscopy Nano-precipitates
We describe here the microstructural evolution, mechanical properties and comprehensive strengthening mechanism of Ti-bearing high strength steels with different finish cooling temperatures in the context of newly developed ultrafast cooling system. Pilot-scale studies demonstrated that high yield strength of ~650 MPa can be obtained with ultrafast cooling finish temperature of 580 °C after hot rolling. The underlying reason is that the microalloyed carbides and nanoscale cementites can precipitate simultaneously to improve the precipitation strengthening to a large extent. In order to estimate the precipitation strengthening, the volume fraction of the precipitates in different size ranges was obtained by chemical phase analysis, small-angle X-ray and neutron scattering. The results indicated that Fe3C, with a higher volume fraction, had a stronger precipitation strengthening effect than nanoscale TiC. The precipitation strengthening contribution of nanoscale precipitates can achieve 350 MPa. Together with solid solution strengthening and grain refinement strengthening, the theoretical calculated values match well with the experimental values.
1. Introduction Low carbon microalloyed steel plates are now widely used in buildings, bridges, ships, vessels, industrial equipment and storage tanks. Excellent mechanical properties with high strength and high ductility can be obtained by optimizing the new generation thermomechanical controlled process (NG-TMCP), which refers to the ultrafast cooling (UFC) combined with thermomechanical controlled processing [1–4]. NG-TMCP has been currently applied in industrial production, aiming at reducing the consumption of alloying elements and making the steel manufacturing process economically viable. The superior mechanical properties are a consequence of grain refinement together with microstructural control and precipitation hardening. The strong carbide-forming elements titanium can not only facilitate grain refinement but also contribute to dispersion hardening through carbide precipitation in the matrix. In addition, UFC can greatly increase the precipitation hardening since it can suppress the amount of precipitates in austenite during the relatively longer rolling time and increase the
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supersaturation ratio in ferrite. Precipitation strengthening has triggered extensive interests since the development of Mo-Ti-bearing steel with a strength of 780 MPa and more recently 1200 MPa through controlled precipitation strengthening. The level of the yield strength increment from precipitation strengthening was almost twice of that obtained by conventional HSLA steels [5,6]. However, steel industries continue to face the challenge of reducing the alloy costs. In order to meet the requirements of cost reduction and maintaining high strength, nanoscale cementite strengthening is regarded as a viable option to reduce the addition of microalloying elements because it is a common secondary phase in steels [7–9]. Wang and his co-workers [7,10,11] investigated the nanoscale cementites in different carbon content steels without alloy addition by controlling the cooling profile after hot rolling. The results indicate that a low UFC finish temperature (600 °C) and 0.17 wt% carbon content are most favorable for nanoscale cementite precipitation. The strengthening effect of nanoscale cementite precipitates can achieve ~158 MPa. Fu [12] proposed that if we can combine micro-
Corresponding author. Corresponding author. E-mail addresses: [email protected] (X.-T. Deng), [email protected] (Z.-D. Wang).
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http://dx.doi.org/10.1016/j.msea.2017.05.066 Received 22 January 2017; Received in revised form 4 May 2017; Accepted 16 May 2017 Available online 17 May 2017 0921-5093/ © 2017 Elsevier B.V. All rights reserved.
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a solution consisting of 650 ml ethyl alcohol, 100 ml perchloric acid, and 50 ml distilled water. Thin foils for transmission electron microscopy (TEM) observations were mechanically thinned to ~0.06 mm and then electrochemically jet polished at −30 °C in a solution containing 9 vol% perchloric acid in ethanol.
alloyed carbide and cementite strengthening together, then the strength will be improved to a large extent. They observed nanoscale cementite and microalloyed carbides precipitates simultaneously in 0.06C-0.10Si0.08Ti-microalloyed high strength weathering steels during thin slab continuous casting and rolling. The nanoscale cementites and microalloyed carbides are both believed to play an important role in enhancing the strength of the steel. The yield strength calculation formula for low carbon steel was also proposed, where the yield strength of steel equals the sum of solid solution strengthening, grain refinement strengthening, and precipitation strengthening [12]. However, little work has been carried out in investigating UFC finish temperature to control nanoscale cementites and microalloyed carbides precipitated simultaneously in low-carbon steels plates during NGTMCP process. In this work, we investigated the effect of UFC finish temperature after hot rolling on the properties of a Ti-bearing low carbon steel plates. Chemical phase analysis, small-angle X-ray and neutron scattering (SAXS and SANS) and high-resolution transmission electron microscopy (HRTEM) were used simultaneously to conduct accurate quantitative characterization of the precipitates, in terms of precipitates size, morphology, chemical composition constituent and volume fraction. In addition, electron back-scattered diffraction (EBSD) and scanning electron microscopy (SEM) were also used to determine the microstructure evolution and fracture morphology, in order to clarify the relationship between microstructure and mechanical properties of the Ti-bearing steel processed by NG-TMCP.
2.3. SANS, chemical phase analysis and SAXS analysis SANS was employed to analyze volume fraction of tiny precipitates. Basics of the SANS technique can be found in literature [13–16]. The scattering intensity I(q) is measured as a function of scattering vector q =4πsinθ/λ, where λ is the wavelength of the incident neutrons, and θ is half of the scattering angle. A monochromatic neutron beam with a mean wavelength of 0.53 nm (Δλ/λ=10%) was produced using a multiblade mechanical velocity selector. The samples were measured under a magnetic field of 1.5 T in order to separate the magnetic and nuclear scattering signals. By changing the sample-detector distances, ranging from 0.08 to 2.0 nm−1 has been covered. In this work, chemical phase analysis and SAXS methods were used simultaneously with TEM and SANS to investigate the chemical composition constituents, volume fraction of nanoscale precipitates with different sizes. The precipitates were electrochemically extracted from the steels, according to the test standard ISO/TS 13762. The detailed procedures for chemical phase analysis were as follows: (1) Electrolytic dissolution of the steel sample to obtain electrolyzed powders containing iron carbides, microalloyed carbides, sulfides and nitrides. (2) Elimination of iron carbides, sulfides, and AlN to obtain microalloyed carbides and oxides. (3) Removal of microalloyed carbides to get stable oxides.
2. Materials and methods 2.1. Materials and thermo-mechanical Processing The nominal composition of the experimental steel in weight percentage (wt%) was 0.15 C, 0.98Mn, 0.28Si, 0.02Al, 0.0027 N, 0.0048 O and balance Fe. A small amount of Ti (0.08 wt%) was added based on the composition Q345 for the maximum precipitation strengthening and grain refinement strengthening. The experimental steels were vacuum-melted in an induction furnace and forged into ingots. After austenitizing at 1250 °C for 2 h, hot rolling was carried out using Φ450 mm experimental hot mill with the finish rolling temperature of 880 °C. Then the steels were ultrafast cooled to different temperatures (620, 580 and 540 °C) and held for 20 min in asbestos. Finally, the steels were air cooled to room temperature. The detailed processing parameters are described in Table 1.
After the above separation procedure, the electrolyzed powders were dissolved and the content of Fe, Mn, N, and C was analyzed. Then the mass fraction of cementite based on formula (FeaMnb)3(CxNy) was calculated. According to GB/T 1322191 (ISO/TS 13762 2001) the size distribution of precipitate powders was analyzed by SAXS with a Kratky small-angle X-ray scattering (Rigaku Corporation, Tokyo, Japan). Using this approach, the analyzed error was less than 10 pct.
2.4. Tensile and impact toughness tests Standard cylindrical tensile test samples with a gage length of 25 mm and diameter of 5 mm were prepared from the steel plates perpendicular to the rolling direction. Tensile tests were conducted at room temperature to measure the yield strength, tensile strength and elongation with a crosshead speed of 1 mm/min, using a SANS-5000 tensile tester. Standard Charpy v-notch impact samples of dimensions 10×10×55 mm3 were prepared along the rolling direction to determine impact toughness at −20 and −40 °C using an Instron 9250 impact tester.
2.2. Microstructural characterization The specimens were cut from the hot rolled plates with their surfaces along thickness direction and rolling direction. The microstructure was observed using a combination of optical microscope (LEICA DMIRM), scanning electron microscope (Quanta 600), and transmission electron microscope (Tecnai G2 F20). For EBSD detections, the specimens were electrochemically polished at room temperature in Table 1 Thermal mechanical processing parameters of the experimental steels. No.
Plate thickness
Rolled in recrystallization temperature region
Rolled in non-recrystallization temperature region
Start, °C
Finish, °C
Start, °C
Finish, °C
Cooling rate, °C/s
Final temperature, °C
Type of cooling
Holding 20 min, air cooling Holding 20 min, air cooling Holding 20 min, air cooling
1
12
1150
1120
882
874
72
620
2
12
1150
1096
889
864
64
580
3
12
1150
1116
880
874
64
540
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Fig. 1. Optical micrographs of steels cooled to (a) 620 °C, (c) 580 °C and (e) 540 °C; representative TEM micrographs of (b) PF with high density dislocation in steel cooled to 620 °C, (d) GB morphology and SAED pattern of M/A constituent in steel cooled to 580 °C and (f) BF in steel cooled to 540 °C. (RD and ND in Fig. 1(a) (c) and (d) represent rolling and normal direction, respectively.).
3. Results and discussion
of steel cooled to 580 °C, which indicates that the microstructure mainly consisted of granular bainite (GB) and a small amount of PF. Fig. 1(d) shows a typical TEM image of GB and the selected area electron diffraction (SAED) pattern of martensite/austenite (M/A) constituent marked by white circle. It demonstrates that some M/A constituents indicated by red arrows are located in ferrite and the martensite obeys the Kurdjumov-Sachs (KS) orientation relationship (OR) with respect to the austinite. The M/A constituents are rich in carbon due to its diffusion during bainitic transformation. In steel cooled to 540 °C, the microstructure primarily consisted of bainite ferrite (BF) and allotriomorphic ferrite (AF), as shown in Fig. 1(e). The
3.1. Microstructural evolution Optical micrographs and TEM micrographs of steels cooled to different temperatures are presented in Fig. 1. The microstructure of steel cooled to 620 °C mainly consisted of polygonal ferrite (PF), as shown in Fig. 1(a) and (b). Typical morphology of high density dislocation in PF can be observed obviously in the TEM micrograph, which was introduced by large reduction of hot rolling and kept by the UFC process after hot rolling. Fig. 1(c) illustrates the optical micrograph
Fig. 2. EBSD orientation maps for steels with UFC finish temperature of (a) 620 °C and (b) 580 °C.
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Fig. 3. TEM images of steel cooled to 620 °C: (a) morphology of TiC and cementite located homogeneously in PF, (b) SAED pattern of the precipitates indicated by blue circle in (a), (c) HRTEM image of the precipitate marked by blue tangle in (a) and (d) the corresponding FFT diffractogram of the marked area as indicated by blue square in (c). (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article).
carbides precipitated in PF in steel cooled to 580 °C. Random distributed microalloyed carbides of size range ~2–10 nm were observed. Representative HRTEM image and the corresponding FFT diffratogram of the nanoscale carbide along the [111] zone axis of ferrite is presented in Fig. 4(b). It follows that these precipitates have a NaCl-type structure and the lattice parameter of the precipitates was calculated to be 0.432 nm. The precipitate with this lattice parameter and NaCl structure has been determined to be TiC. From the FFT diffratogram, it can be concluded that the TiC carbide obeys the OR of [111]ferrite //[110]TiC and (101)ferrite //(002)TiC with respect to the ferrite matrix. Fig. 4(c) shows a dark field (DF) image obtained by using the (210) reflection in the SAED pattern along [001] zone axis of the other kind of precipitates formed in steel cooled to 580 °C. The precipitates are in the range of ~10–30 nm and can be confirmed to be cementites by SAED pattern shown in the inset in Fig. 4(c). Fig. 5(a) shows the TEM image of the steel cooled to 540 °C and the SAED pattern of cementite along its [210] zone axis. It is demonstrated that the microstructure mainly consisted of bainite laths and some flake-shaped cementites (as pointed by white arrows) were in the boundaries between bainite laths. The average length and width of these cementites was determined to be ~400 nm and ~50 nm, respectively. Fig. 5(b) shows the TEM image of the nanoscale precipitates formed in the bainite laths and the corresponding SEAD pattern along its [011] zone axis. It was confirmed to be TiC. Representative HRTEM images of nanometer-sized TiC and Fe3C are presented in Fig. 6. As illustrated in Fig. 6(a), TiC is very fine and within the thickness of foil. This leads to the development of Moiré fringe contrast because of overlapping of carbide and ferrite lattice, which provides a clear contrast of the carbide and can be used for the accurate
hot rolling deformation and deformation induced precipitates in austenite can accelerate the nucleation of the AF on the austenite grain boundary. Fig. 1(f) shows the TEM image of BF with some fine flaketype carbides located between the BF laths. Fig. 2 shows the orientation maps of steels cooled to 620 and 580 °C, analyzed by EBSD. In the orientation image map of steel cooled to 620 °C (Fig. 2(a)), it is clear that the grains have obvious preferred orientations, with strong < 101 > orientations, which is not observed in steel cooled to 580 °C, as shown in Fig. 2(b). The effective grain size was evaluated by equivalent grain boundaries of larger than 15° per unit area. In this study, the average effective grain size is 7.5 µm and 6.9 µm in steels cooled to 620 and 580 °C, respectively.
3.2. Precipitation behavior Fig. 3(a) shows TEM image of microalloyed carbides and cementites precipitated homogeneously in PF in steel cooled to 620 °C. Fig. 3(b) illustrates the SAED pattern of the comparatively large precipitates marked by blue circle in Fig. 3(a) along the [301] zone axis. It indicates that these precipitates have an orthogonal structure. The lattice parameters a and b of the precipitate was calculated to be 0.449 nm and 0.506 nm, which is in accordance with that of cementite. Fig. 3(c) shows the HRTEM image of the precipitate marked by blue square in Fig. 3(a). Fig. 3(d) illustrates the corresponding fast Fourier transformation (FFT) diffratogram. The lattice parameter of the precipitates was calculated to be 0.432 nm, coincident with that of TiC. In addition, the TiC carbides can be determined to obey the OR of [111]ferrite //[110]TiC and (101)ferrite //(002)TiC with respect to the ferrite matrix. Fig. 4(a) shows the morphology and distribution of the nanoscale 271
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Fig. 4. TEM images of steel cooled to 580 °C: (a) morphology of TiC in PF, (b) HRTEM image of TiC particle and the corresponding FFT diffratogram, (c) DF image of Fe3C and the corresponding SAED pattern.
at the UFC finish temperature of 580 °C. It gives yield strength 650 MPa, tensile strength 750 MPa, elongation 17.4% and impact energy at test temperature of −20 °C 93.4 J. The introduction of only 0.08 wt% Ti combined with NG-TMCP demonstrates impressive improvement of the strength compared with Q345 steels. However, the yield strength and impact energy (−20 °C) were 590 MPa and 26.3 J for the steel cooled to 540 °C. The poor mechanical properties can be attributed to the large size cementites formed on the bainite lath boundaries. According to the Griffth crack propagation theory, the coarse hard phase assists microcrack initiation by reducing the initiation energy [17]. In addition, the cementites formed on the lath boundaries reduce the potential of nanoscale cementite strengthening contribution to yield strength. So, the following characterization mainly focuses on steel cooled to 620 and 580 °C. The SEM fractographs of experimental steels after impact test at −40 °C are presented in Fig. 8. The macrograph fracture morphology of steel cooled to 620 °C is shown in Fig. 8(a). The zone 1 in the macrograph stands for radical zone. The fractograph of radical zone was predominantly cleavage fracture with some obvious cracks indicated by red arrows, as shown in Fig. 8(b). Fig. 8(c) shows the macrograph fracture morphology of steel cooled to 580 °C. The zone 1 is radical zone and the zone 2 is shear lip zone. The fractograph of radical zone was predominantly quasi-cleavage facets with some small deep dimples (marked by blue arrows) and small microcracks (marked by red arrows), as shown in Fig. 8(d). Fig. 8(e) shows the fractograph of shear lip zone, which mainly consisted of small deep dimples with small precipitates (marked by pink circles) in it. In steel cooled to 620 °C,
measurement of the carbide size. The size of the TiC in Fig. 6(a) was determined to be ~4.41 nm along the direction of the Moiré fringe (length) and ~4.32 nm perpendicular to the direction of the Moiré fringe (thickness). The aspect ratio (length/thickness) is close to 1. Thus, it can be concluded that the morphology of the carbide was close to spherical. The average value of length and thickness of TiC was ~4.36 nm. Futhermore, Fig. 6(a) also reveals an excellent coherent relationship between the TiC phase and the matix. The nanoscale coherent interfaces between TiC and the matrix have a low interaface energy and good microsturcture stablity. Fig. 6(b) shows the representive HRTEM image of the Fe3C, the lattice parameter can be calculated and the results match well with the SAED pattern. Fig. 7 shows the size distribution histograms of TiC and cementite precipitated in steel cooled to 620 and 580 °C. With each speicmen at least 100 precipitates was measured. The averge sizes of TiC and cementite was 5.4 and 32.8 nm in steel cooled to 620 °C, while for the steel cooled to 580 °C, the averge size of TiC and cementite was 3.8 and 22.4 nm. 3.3. Mechanical properties Table 2 summarizes the mechanical properties of the experimental steels after hot rolling with different UFC finish temperatures, including yield strength, tensile strength, total elongation and Charpy impact energy. Each datum is obtained from three specimens. The tensile and Charpy impact test results indicate that the mechanical property reaches optimal with excellent combination of strength and toughness
Fig. 5. TEM images of steel cooled to 540 °C: (a) the cementites located on the boundaries of bainite laths and (b) the TiC precipitate formed in bainite lath; the insets in (a) and (b) are the SEAD patterns of cementite and TiC.
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Fig. 6. Representative HRTEM images of (a) TiC and (b) cementite particles.
Fig. 7. Size distribution histograms of TiC and cementite precipitated in steel cooled to (a) 580 °C and (b) 620 °C.
3.4. Strengthening mechanism
Table 2 Mechanical properties with standard deviation of the experimental steels with different UFC finish temperatures. Steel
1 2 3
Yield strength, MPa
615 (2.8) 650 (4.3) 590 (2.1)
Tensile strength, MPa
735 (5.7) 750 (2.8) 730 (6.1)
Elongation, %
22.6 (1.3) 17.4 (1.1) 10.9 (0.9)
Chemical phase analysis and SAXS analysis were carried out on steels cooled to 620 °C and 580 °C. Table 3 summarizes the structural parameters of the extracted precipitates by electrolysis using X-ray diffraction. It can be concluded from the XRD results that the precipitates were mainly composed of M3(C, N), Ti2CS, TiC and Ti(C, N). The weight fraction of MC and M3C are presented in Table 4. The content of M3C cementite-based precipitates dominated compared with that of MC. The weight fraction of M3C and MC type precipitates is similar in steel cooled 620 °C and 580 °C. The volume fraction and the corresponding strengthening increment of M3C and MC type precipitates in different diameter ranges below ~ 36 nm are shown in Table 5. It can be noted that the volume fraction of M3C and MC carbides with size less than ~36 nm in steel cooled to 620 °C is smaller than that in steel cooled to 580 °C, which will result in difference in strengthening contribution. The main objective of NG-TMCP is to enhance the strength of steel by precipitation hardening. It is therefore essential to estimate the contribution of the precipitates toward effective strengthening. There are two well-known mechanisms for precipitation strengthening, shearing mechanism and bypass mechanism [1]. The equations for the two contribution are as followings [23].
Impact toughness (V-notch, J) −20 °C
−40 °C
62.4 (7.2) 93.4 (8.6) 26.3 (3.8)
43.5 (3.6) 65.7 (6.7) 18.4 (2.4)
seldom shear lip zone can be observed, indicative of poor lowtemperature toughness compared with steel cooled to 580 °C. The main reason for the low impact energy in steel cooled to 620 °C is the existence of relatively large size cementites compared with that in steel cooled to 580 °C. Another factor that affects low temperature toughness should be related to microstructure. Fig. 9(a) and (b) shows the quality maps of steels cooled to 620 and 580 °C, respectively. The red lines represent the small angle grain boundaries with misorientation of 2–15°, while the black lines correspond to high angle grain boundaries with misoientation of 15° and larger than 15°. The high misorientation grain/packet boundaries can effectively deflect or even arrest the propagation of cleavage microcracks, whereas the low misorientation boundaries have less ability to deflect the crack [18,19]. It is obvious to notice that the percentage of high misoriention boundary in steel cooled to 580 °C is higher compared that in steel cooled to 620 °C, as shown in Fig. 9(c) and (d). These high grain boundaries in steel cooled to 580 °C contribute more resistance to cleavage microcrack [20–22].
σbypass = 10.8
f d
σshearing = M τP = 273
ln(1630d )
γ 3/2 2 × 1.1 × 2 × d1/2f 1/2 b 2AG
(1)
(2)
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Fig. 8. The impact fractographs: (a) macrograph and (b) the fractograph of zone 1 (radical zone) in steel cooled to 620 °C; (c) macrograph and the fractographs of (d) zone 1 (radical zone) and (e) zone 2 (shear lip zone) in steel cooled to 620 °C.
Fig. 9. EBSD image quality maps with grain boundary misorientation distribution of steels with UFC finish cooling temperature of (a) 620 °C and (b) 580 °C; grain boundary distribution histograms of steels with UFC finish cooling temperature of (c) 620 °C and (d) 580 °C.
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Table 3 Structural parameters of the precipitates.
Table 6 Comparison of calculated yield strength values with experimental values in steels with UFC finish temperatures of 620 and 580 °C.
Phase structure
Lattice constant, nm
M3(C,N)
a0=0.4523–0.4530, b0=0.5088–0.5080, c0=0.6743–0.6772 a0=0.3210–0.3240， c0=01.1203–1.1308， c/a=3.49 a0=0.431–0.433 a0=0.425–0.427
Ti2CS TiC Ti(C,N)
Crystal system No.
D, μm
Orthorhombic 1 2
Hexagonal
7.5 6.9
Calculated Yield Strength Values, MPa σGS
σSS
σSP
σy
*σSP
*σy
Actual measured σy, MPa
218.9 228.3
77.7 78.6
279.4 306.9
576.0 613.8
314.4 351.6
611.0 658.5
615 650
Face centered cubic Face centered cubic
Table 4 The mass fraction and chemical structural formula in steel with UFC finish temperature of 620 and 580 °C. No.
1 2
MC phase
M3C phase
Mass fraction (%)
Structural formula
Mass fraction (%)
Structural formula
0.0313 0.0338
Ti(C0.609N0.391) Ti(C0.672N0.328)
1.4113 1.5553
(Fe0.9856Mn0.0144)3C (Fe0.984Mn0.016)3C
Where, τP represents the shear stress caused by dislocations cutting the particles in MPa; b is dislocation Burger's vector and equal to 0.248 nm; G is the shear elasticity modulus and equal to 80650 MPa; γ is the interface energy between precipitates and matrix and equal to 0.5–1 J/ m2; d represents the second-phase particle diameter in μm; f represents the volume fraction of the precipitates; From formula (1) and (2), we can notice that the strengthening increment of shearing mechanism increases with the increasing precipitate size, while the strengthening increment of the bypass mechanism decreases with the increase of particle size. The critical transformation size dc can be calculated by the following equation.
Gb 2 d ln( c ) dc = 0.209 Kγ 2b
Fig. 10. The magnetic scattering of steels cooled to 620 and 580 °C by SANS. (Measured data are presented in solid square or circle, fitted data are presented in solid line).
σy = σSG + σSS + σSP = 600D−1/2 + {46[C ] + 37[Mn] + 83[Si] + 59[Al] + 2918[N ] + 80.5[Ti]} + σSP
(4)
where σy , σSG , σSS and σSP represent the yield strength, grain refinement strengthening, solid solution strengthening, and precipitation strengthening in MPa, respectively. The components of the yield strength and the total yield strength are presented in Table 6. The result shows that estimated yield strength for steels cooled to 620 and 580 °C are 576.0 MPa and 613.8 MPa, lower than the actual yield strength. Exploring the reason, it should be attributed to the limitation of electrolytic separation that some fine precipitates lost during the electrolyzed powders filter process. So, the SANS was chosen to confirm the volume fraction of tiny precipitates with bulk samples. Fig. 10 shows the magnetic scattering curves of steels cooled to 620 °C and 580 °C. The magnetic scattering intensity was extracted from the intensity collected when the scattering vector is perpendicular (α=90°) and parallel (α=0°) to the applied magnetic field, according to
(3)
The meaning of symbols in Eq. (3) is same as that mentioned above. The critical transformation size of the precipitate depends on the interface energy between the precipitates and the matrix. The dc of TiC and Fe3C were estimated to be ~1.5–6 nm and ~4.7–10 nm calculated by Eq. (3). The precipitation strengthening effect of TiC in all size ranges and Fe3C with size large than 10 nm were calculated based on the bypass mechanism. For Fe3C with size less than 10 nm, the shearing mechanism was applied [12]. The calculated results of precipitation strengthening increment in different size range are listed in Table 6. For low-carbon steel, the yield strength equals to the sum of solid strengthening, grain refinement strengthening, and precipitation strengthening, and is given by the following equation [12].
Imagnetic = I (α = 90°) − I (α = 0°)
(5)
Table 5 Contribution of Fe3C and TiC to the yield strength of steels with UFC finish temperatures of 620 and 580 °C. No.
Diameter Range, nm
Fe3C
TiC
Total increment, Mpa
Volume Fraction, Pct
Yield Strength Increment, MPa
Volume Fraction, Pct
Yield strength increment, MPa
1
1–5 5–10 10–18 18–36 ∑
0.0153 0. 0038 0.0020 0.0025 0.0236
60.5 16.5 7.6 5.2 89.7
0. 0364 0 0. 0873 0. 3550 0.4786
78.6 0.0 50.1 61.0 189.7
279.4
2
1–5 5–10 10–18 18–36 ∑
0.0060 0.0026 0.0011 0.0032 0.0130
38.0 13.6 5.6 5.9 63.1
0.1172 0.0035 0.0019 0.2226 0.3452
153.9 34.2 7.4 48.3 243.8
306.9
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3. The contribution of precipitates to yield strength was greater than ~350 MPa steel with UFC finish cooling temperature, which clearly indicates that precipitation strengthening primarily contributed to the total yield strength. A small amount of Ti (0.08 wt%) addition can improve the traditional Q345 steel to obtain 650 MPa in yield strength.
Table 7 The fine precipitate size and volume fraction in steels with UFC finish temperatures of 620 and 580 °C measured by SANS. No.
1 2
SANS Rs (nm)
Gs
f
2.4(0.2) 1.8(0.1)
0.09(0.02) 0.17(0.03)
0.055% 0.130%
Acknowledgements The research was supported financially by National Science Foundation of China (Grant No. 51234002, 51504064 and 51474064) and National Key Research and Development Program 2016YFB0300601. The China Scholarship Council is also acknowledged sincerely for supporting us to complete this work. RDKM gratefully acknowledges support from University of Texas at El Paso, USA.
Three different slopes can be observed for both samples. The low-q data decays with q−4, which is induced by grain boundaries and large precipitates. The slope of the median- and high-q data is −2.7 and −1.8 in logarithmic coordinate, respectively. It means that the precipitates have two characteristic sizes. We fit the data with a power law and Guinier laws due to the limited fitting range,
I (q ) = Gs exp(−
q 2Rg2 q 2Rs2 ) + G exp(− ) + Aq−4 3 3
References (6)
[1] T. Gladman, The Physical Metallurgy of Microalloyed Steels, The Institute of Materials, London, 1997. [2] L. Cheng, Q.W. Cai, B.S. Xie, Z. Ning, X.C. Zhou, G.S. Li, Mater. Sci. Eng. A651 (2016) 185–191. [3] S. Tang, Z.Y. Liu, G.D. Wang, R.D.K. Misra, Mater. Sci. Eng. A 580 (2013) 257–265. [4] X.H. Cai, C.B. Liu, Z.Y. Liu, Mater. Des. 53 (2014) 998–1004. [5] N. Kamikawa, Y. Abe, G. Miyamoto, Y. Funakawa, T. Furuhara, ISIJ Int. 54 (2014) 212–221. [6] Y. Funakawa, T. Shiozaki, K. Tomita, T. Yamamoto, E. Maeda, ISIJ Int. 44 (2004) 1945–1951. [7] B. Wang, Z.Y. Liu, X.G. Zhou, G.D. Wang, R.D.K. Misra, Mater. Sci. Eng. A 575 (2013) 189–198. [8] M.C. Zhao, X.F. Huang, A. Atrens, Mater. Sci. Eng. A 549 (2012) 222–227. [9] Q.L. Yong, Secondary Phase in Steels, Metallurgical Industry Press, Beijing, 2006. [10] J. Fu, G.Q. Li, X.P. Mao, K.M. Fang, Metall. Mater. Trans. A 42 (2011) 3797–3812. [11] B. Wang, Z.D. Wang, B.X. Wang, G.D. Wang, R.D.K. Misra, Metall. Mater. Trans. A 46 (2015) 2834–2843. [12] B. Wang, Z.Y. Liu, X.G. Zhou, G.D. Wang, R.D.K. Misra, Mater. Sci. Eng. A 588 (2013) 167–174. [13] B.S. Seong, E. Shin, S.H. Choi, Y. Choi, Y.S. Han, K.H. Lee, Y. Tomota, Appl. Phys. A 99 (2010) 613–620. [14] H. Leitner, P. Staron, H. Clemens, S. Marsoner, P. Warbichler, Mater. Sci. Eng. A 398 (2005) 323–331. [15] P. Staron, B. Jamnig, H. Leitner, R. Ebner, H. Clemensa, J. Appl. Cryst. 36 (2003) 415–419. [16] M. Peng, L.W. Sun, L. Chen, G.A. Sun, B. Chen, C.M. Xie, Q.Z. Xia, G.Y. Yan, Q. Tian, C.Q. Huang, B.B. Pang, Y. Zhang, Y. Wang, Y.G. Liu, W. Kang, J. Gong, Nucl. Instrum. Methods Phys. Res. Sect. A 810 (2016) 63–67. [17] S.H. Goods, L.M. Brown, Acta Metall. 27 (1979) 1–15. [18] J. Hu, L.X. Du, J.J. Wang, C.R. Gao, Mater. Sci. Eng. A 577 (2013) 161–168. [19] X.L. Wan, K.M. Wu, G. Huang, K.C. Nune, Y. Li, L. Cheng, Sci. Technol. Weld. Join. 21 (2016) 295–302. [20] J. Hu, L.-X. Du, J.-J. Wang, H. Xie, C.-R. Gao, R.D.K. Misra, Mater. Sci. Eng. A 585 (2013) 197–204. [21] Y.M. Kim, H. Lee, N.J. Kim, Mater. Sci. Eng. A 478 (2008) 361–370. [22] J. Hu, L.X. Du, J.J. Wang, Q.Y. Sun, Mater. Des. 53 (2014) 332–337. [23] Q.L. Yong, M.T. Ma, B.R. Wu, Microalloyed Steel-Physical and Mechanical Metallurgy, Mechanical Industry Press, Beijing, 1989.
where, Rs and Rg are the size-weighted Guinier radiuses of precipitates with small and medium size, GS and G are the scale factors. The last term in Eq. (6) accounts for the scattering from large precipitates and grain boundaries. The fitting parameters are listed in Table 7. The average radius and the volume fraction of the tiny precipitates (1–5 nm) in steel cooled to 620 °C are 2.4 nm and 0.055%. For steel cooled to 580 °C, they are 1.8 nm and 0.130%, respectively. The precipitation strengthening increment in precipitates (1–5 nm) was estimated to be 174.1 and 198.6 MPa in steels cooled to 620 and 580 °C, respectively. The total precipitation strengthening (*σSP) for the two steels were 314.4 and 351.6 MPa, respectively. The calculated yield strength (*σy) matches well with the measured yield strength, as shown in Table 6. It can be concluded that the contribution of precipitates to yield strength can be greater than ~350 MPa in steel cooled to 580 °C, which clearly indicates that precipitation strengthening primarily contributed to the total yield strength. 4. Conclusion 1. Besides nanoscale TiC, cementite precipitates of size less than ~36 nm were also observed in Ti-microalloyed steel. The volume fraction of Fe3C was significantly higher than TiC for similar size range. 2. When the steel was ultrafast cooled to 580 °C, combination of high yield strength of 650 MPa, good elongation of 17%, and excellent toughness of 90 J at −40 °C was obtained through nanoscale cementites and microalloyed TiC carbides precipitation strengthening.
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Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Nanoscale cementite and microalloyed carbide strengthened Ti bearing low carbon steel plates in the context of newly developed ultrafast cooling
MARK
⁎
X.-L. Lia, C.-S. Leia, Q. Tianb, X.-T. Denga, , L. Chenb, P.-L. Gaob, K.-P. Duc, Y. Dud, Y.-G. Yuc, ⁎⁎ Z.-D. Wanga, , R.D.K. Misrae a
State Key Laboratory of Rolling and Automation, Northeastern, Shenyang 110819, China Key Laboratory of Neutron Physics and Institute of Nuclear Physics and Chemistry, China Academy of Engineering Physics, Mianyang 621999, China c Beijing General Research Institute of Mining and Metallurgy, Beijing 100160, China d Materials Science and Engineering, Northwestern University, Evanston 60208, USA e Laboratory for Excellence in Advanced Steel Research, Department of Metallurgical, Materials and Biomedical Engineering, University of Texas at El Paso, El Paso, TX 79968-0521, USA b
A R T I C L E I N F O
A B S T R A C T
Keywords: High strength low alloy Thermomechanical controlled processing Small-angle neutron and X-ray scattering Microstructural evolution High resolution transmission electron microscopy Nano-precipitates
We describe here the microstructural evolution, mechanical properties and comprehensive strengthening mechanism of Ti-bearing high strength steels with different finish cooling temperatures in the context of newly developed ultrafast cooling system. Pilot-scale studies demonstrated that high yield strength of ~650 MPa can be obtained with ultrafast cooling finish temperature of 580 °C after hot rolling. The underlying reason is that the microalloyed carbides and nanoscale cementites can precipitate simultaneously to improve the precipitation strengthening to a large extent. In order to estimate the precipitation strengthening, the volume fraction of the precipitates in different size ranges was obtained by chemical phase analysis, small-angle X-ray and neutron scattering. The results indicated that Fe3C, with a higher volume fraction, had a stronger precipitation strengthening effect than nanoscale TiC. The precipitation strengthening contribution of nanoscale precipitates can achieve 350 MPa. Together with solid solution strengthening and grain refinement strengthening, the theoretical calculated values match well with the experimental values.
1. Introduction Low carbon microalloyed steel plates are now widely used in buildings, bridges, ships, vessels, industrial equipment and storage tanks. Excellent mechanical properties with high strength and high ductility can be obtained by optimizing the new generation thermomechanical controlled process (NG-TMCP), which refers to the ultrafast cooling (UFC) combined with thermomechanical controlled processing [1–4]. NG-TMCP has been currently applied in industrial production, aiming at reducing the consumption of alloying elements and making the steel manufacturing process economically viable. The superior mechanical properties are a consequence of grain refinement together with microstructural control and precipitation hardening. The strong carbide-forming elements titanium can not only facilitate grain refinement but also contribute to dispersion hardening through carbide precipitation in the matrix. In addition, UFC can greatly increase the precipitation hardening since it can suppress the amount of precipitates in austenite during the relatively longer rolling time and increase the
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supersaturation ratio in ferrite. Precipitation strengthening has triggered extensive interests since the development of Mo-Ti-bearing steel with a strength of 780 MPa and more recently 1200 MPa through controlled precipitation strengthening. The level of the yield strength increment from precipitation strengthening was almost twice of that obtained by conventional HSLA steels [5,6]. However, steel industries continue to face the challenge of reducing the alloy costs. In order to meet the requirements of cost reduction and maintaining high strength, nanoscale cementite strengthening is regarded as a viable option to reduce the addition of microalloying elements because it is a common secondary phase in steels [7–9]. Wang and his co-workers [7,10,11] investigated the nanoscale cementites in different carbon content steels without alloy addition by controlling the cooling profile after hot rolling. The results indicate that a low UFC finish temperature (600 °C) and 0.17 wt% carbon content are most favorable for nanoscale cementite precipitation. The strengthening effect of nanoscale cementite precipitates can achieve ~158 MPa. Fu [12] proposed that if we can combine micro-
Corresponding author. Corresponding author. E-mail addresses: [email protected] (X.-T. Deng), [email protected] (Z.-D. Wang).
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http://dx.doi.org/10.1016/j.msea.2017.05.066 Received 22 January 2017; Received in revised form 4 May 2017; Accepted 16 May 2017 Available online 17 May 2017 0921-5093/ © 2017 Elsevier B.V. All rights reserved.
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a solution consisting of 650 ml ethyl alcohol, 100 ml perchloric acid, and 50 ml distilled water. Thin foils for transmission electron microscopy (TEM) observations were mechanically thinned to ~0.06 mm and then electrochemically jet polished at −30 °C in a solution containing 9 vol% perchloric acid in ethanol.
alloyed carbide and cementite strengthening together, then the strength will be improved to a large extent. They observed nanoscale cementite and microalloyed carbides precipitates simultaneously in 0.06C-0.10Si0.08Ti-microalloyed high strength weathering steels during thin slab continuous casting and rolling. The nanoscale cementites and microalloyed carbides are both believed to play an important role in enhancing the strength of the steel. The yield strength calculation formula for low carbon steel was also proposed, where the yield strength of steel equals the sum of solid solution strengthening, grain refinement strengthening, and precipitation strengthening [12]. However, little work has been carried out in investigating UFC finish temperature to control nanoscale cementites and microalloyed carbides precipitated simultaneously in low-carbon steels plates during NGTMCP process. In this work, we investigated the effect of UFC finish temperature after hot rolling on the properties of a Ti-bearing low carbon steel plates. Chemical phase analysis, small-angle X-ray and neutron scattering (SAXS and SANS) and high-resolution transmission electron microscopy (HRTEM) were used simultaneously to conduct accurate quantitative characterization of the precipitates, in terms of precipitates size, morphology, chemical composition constituent and volume fraction. In addition, electron back-scattered diffraction (EBSD) and scanning electron microscopy (SEM) were also used to determine the microstructure evolution and fracture morphology, in order to clarify the relationship between microstructure and mechanical properties of the Ti-bearing steel processed by NG-TMCP.
2.3. SANS, chemical phase analysis and SAXS analysis SANS was employed to analyze volume fraction of tiny precipitates. Basics of the SANS technique can be found in literature [13–16]. The scattering intensity I(q) is measured as a function of scattering vector q =4πsinθ/λ, where λ is the wavelength of the incident neutrons, and θ is half of the scattering angle. A monochromatic neutron beam with a mean wavelength of 0.53 nm (Δλ/λ=10%) was produced using a multiblade mechanical velocity selector. The samples were measured under a magnetic field of 1.5 T in order to separate the magnetic and nuclear scattering signals. By changing the sample-detector distances, ranging from 0.08 to 2.0 nm−1 has been covered. In this work, chemical phase analysis and SAXS methods were used simultaneously with TEM and SANS to investigate the chemical composition constituents, volume fraction of nanoscale precipitates with different sizes. The precipitates were electrochemically extracted from the steels, according to the test standard ISO/TS 13762. The detailed procedures for chemical phase analysis were as follows: (1) Electrolytic dissolution of the steel sample to obtain electrolyzed powders containing iron carbides, microalloyed carbides, sulfides and nitrides. (2) Elimination of iron carbides, sulfides, and AlN to obtain microalloyed carbides and oxides. (3) Removal of microalloyed carbides to get stable oxides.
2. Materials and methods 2.1. Materials and thermo-mechanical Processing The nominal composition of the experimental steel in weight percentage (wt%) was 0.15 C, 0.98Mn, 0.28Si, 0.02Al, 0.0027 N, 0.0048 O and balance Fe. A small amount of Ti (0.08 wt%) was added based on the composition Q345 for the maximum precipitation strengthening and grain refinement strengthening. The experimental steels were vacuum-melted in an induction furnace and forged into ingots. After austenitizing at 1250 °C for 2 h, hot rolling was carried out using Φ450 mm experimental hot mill with the finish rolling temperature of 880 °C. Then the steels were ultrafast cooled to different temperatures (620, 580 and 540 °C) and held for 20 min in asbestos. Finally, the steels were air cooled to room temperature. The detailed processing parameters are described in Table 1.
After the above separation procedure, the electrolyzed powders were dissolved and the content of Fe, Mn, N, and C was analyzed. Then the mass fraction of cementite based on formula (FeaMnb)3(CxNy) was calculated. According to GB/T 1322191 (ISO/TS 13762 2001) the size distribution of precipitate powders was analyzed by SAXS with a Kratky small-angle X-ray scattering (Rigaku Corporation, Tokyo, Japan). Using this approach, the analyzed error was less than 10 pct.
2.4. Tensile and impact toughness tests Standard cylindrical tensile test samples with a gage length of 25 mm and diameter of 5 mm were prepared from the steel plates perpendicular to the rolling direction. Tensile tests were conducted at room temperature to measure the yield strength, tensile strength and elongation with a crosshead speed of 1 mm/min, using a SANS-5000 tensile tester. Standard Charpy v-notch impact samples of dimensions 10×10×55 mm3 were prepared along the rolling direction to determine impact toughness at −20 and −40 °C using an Instron 9250 impact tester.
2.2. Microstructural characterization The specimens were cut from the hot rolled plates with their surfaces along thickness direction and rolling direction. The microstructure was observed using a combination of optical microscope (LEICA DMIRM), scanning electron microscope (Quanta 600), and transmission electron microscope (Tecnai G2 F20). For EBSD detections, the specimens were electrochemically polished at room temperature in Table 1 Thermal mechanical processing parameters of the experimental steels. No.
Plate thickness
Rolled in recrystallization temperature region
Rolled in non-recrystallization temperature region
Start, °C
Finish, °C
Start, °C
Finish, °C
Cooling rate, °C/s
Final temperature, °C
Type of cooling
Holding 20 min, air cooling Holding 20 min, air cooling Holding 20 min, air cooling
1
12
1150
1120
882
874
72
620
2
12
1150
1096
889
864
64
580
3
12
1150
1116
880
874
64
540
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Fig. 1. Optical micrographs of steels cooled to (a) 620 °C, (c) 580 °C and (e) 540 °C; representative TEM micrographs of (b) PF with high density dislocation in steel cooled to 620 °C, (d) GB morphology and SAED pattern of M/A constituent in steel cooled to 580 °C and (f) BF in steel cooled to 540 °C. (RD and ND in Fig. 1(a) (c) and (d) represent rolling and normal direction, respectively.).
3. Results and discussion
of steel cooled to 580 °C, which indicates that the microstructure mainly consisted of granular bainite (GB) and a small amount of PF. Fig. 1(d) shows a typical TEM image of GB and the selected area electron diffraction (SAED) pattern of martensite/austenite (M/A) constituent marked by white circle. It demonstrates that some M/A constituents indicated by red arrows are located in ferrite and the martensite obeys the Kurdjumov-Sachs (KS) orientation relationship (OR) with respect to the austinite. The M/A constituents are rich in carbon due to its diffusion during bainitic transformation. In steel cooled to 540 °C, the microstructure primarily consisted of bainite ferrite (BF) and allotriomorphic ferrite (AF), as shown in Fig. 1(e). The
3.1. Microstructural evolution Optical micrographs and TEM micrographs of steels cooled to different temperatures are presented in Fig. 1. The microstructure of steel cooled to 620 °C mainly consisted of polygonal ferrite (PF), as shown in Fig. 1(a) and (b). Typical morphology of high density dislocation in PF can be observed obviously in the TEM micrograph, which was introduced by large reduction of hot rolling and kept by the UFC process after hot rolling. Fig. 1(c) illustrates the optical micrograph
Fig. 2. EBSD orientation maps for steels with UFC finish temperature of (a) 620 °C and (b) 580 °C.
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Fig. 3. TEM images of steel cooled to 620 °C: (a) morphology of TiC and cementite located homogeneously in PF, (b) SAED pattern of the precipitates indicated by blue circle in (a), (c) HRTEM image of the precipitate marked by blue tangle in (a) and (d) the corresponding FFT diffractogram of the marked area as indicated by blue square in (c). (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article).
carbides precipitated in PF in steel cooled to 580 °C. Random distributed microalloyed carbides of size range ~2–10 nm were observed. Representative HRTEM image and the corresponding FFT diffratogram of the nanoscale carbide along the [111] zone axis of ferrite is presented in Fig. 4(b). It follows that these precipitates have a NaCl-type structure and the lattice parameter of the precipitates was calculated to be 0.432 nm. The precipitate with this lattice parameter and NaCl structure has been determined to be TiC. From the FFT diffratogram, it can be concluded that the TiC carbide obeys the OR of [111]ferrite //[110]TiC and (101)ferrite //(002)TiC with respect to the ferrite matrix. Fig. 4(c) shows a dark field (DF) image obtained by using the (210) reflection in the SAED pattern along [001] zone axis of the other kind of precipitates formed in steel cooled to 580 °C. The precipitates are in the range of ~10–30 nm and can be confirmed to be cementites by SAED pattern shown in the inset in Fig. 4(c). Fig. 5(a) shows the TEM image of the steel cooled to 540 °C and the SAED pattern of cementite along its [210] zone axis. It is demonstrated that the microstructure mainly consisted of bainite laths and some flake-shaped cementites (as pointed by white arrows) were in the boundaries between bainite laths. The average length and width of these cementites was determined to be ~400 nm and ~50 nm, respectively. Fig. 5(b) shows the TEM image of the nanoscale precipitates formed in the bainite laths and the corresponding SEAD pattern along its [011] zone axis. It was confirmed to be TiC. Representative HRTEM images of nanometer-sized TiC and Fe3C are presented in Fig. 6. As illustrated in Fig. 6(a), TiC is very fine and within the thickness of foil. This leads to the development of Moiré fringe contrast because of overlapping of carbide and ferrite lattice, which provides a clear contrast of the carbide and can be used for the accurate
hot rolling deformation and deformation induced precipitates in austenite can accelerate the nucleation of the AF on the austenite grain boundary. Fig. 1(f) shows the TEM image of BF with some fine flaketype carbides located between the BF laths. Fig. 2 shows the orientation maps of steels cooled to 620 and 580 °C, analyzed by EBSD. In the orientation image map of steel cooled to 620 °C (Fig. 2(a)), it is clear that the grains have obvious preferred orientations, with strong < 101 > orientations, which is not observed in steel cooled to 580 °C, as shown in Fig. 2(b). The effective grain size was evaluated by equivalent grain boundaries of larger than 15° per unit area. In this study, the average effective grain size is 7.5 µm and 6.9 µm in steels cooled to 620 and 580 °C, respectively.
3.2. Precipitation behavior Fig. 3(a) shows TEM image of microalloyed carbides and cementites precipitated homogeneously in PF in steel cooled to 620 °C. Fig. 3(b) illustrates the SAED pattern of the comparatively large precipitates marked by blue circle in Fig. 3(a) along the [301] zone axis. It indicates that these precipitates have an orthogonal structure. The lattice parameters a and b of the precipitate was calculated to be 0.449 nm and 0.506 nm, which is in accordance with that of cementite. Fig. 3(c) shows the HRTEM image of the precipitate marked by blue square in Fig. 3(a). Fig. 3(d) illustrates the corresponding fast Fourier transformation (FFT) diffratogram. The lattice parameter of the precipitates was calculated to be 0.432 nm, coincident with that of TiC. In addition, the TiC carbides can be determined to obey the OR of [111]ferrite //[110]TiC and (101)ferrite //(002)TiC with respect to the ferrite matrix. Fig. 4(a) shows the morphology and distribution of the nanoscale 271
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Fig. 4. TEM images of steel cooled to 580 °C: (a) morphology of TiC in PF, (b) HRTEM image of TiC particle and the corresponding FFT diffratogram, (c) DF image of Fe3C and the corresponding SAED pattern.
at the UFC finish temperature of 580 °C. It gives yield strength 650 MPa, tensile strength 750 MPa, elongation 17.4% and impact energy at test temperature of −20 °C 93.4 J. The introduction of only 0.08 wt% Ti combined with NG-TMCP demonstrates impressive improvement of the strength compared with Q345 steels. However, the yield strength and impact energy (−20 °C) were 590 MPa and 26.3 J for the steel cooled to 540 °C. The poor mechanical properties can be attributed to the large size cementites formed on the bainite lath boundaries. According to the Griffth crack propagation theory, the coarse hard phase assists microcrack initiation by reducing the initiation energy [17]. In addition, the cementites formed on the lath boundaries reduce the potential of nanoscale cementite strengthening contribution to yield strength. So, the following characterization mainly focuses on steel cooled to 620 and 580 °C. The SEM fractographs of experimental steels after impact test at −40 °C are presented in Fig. 8. The macrograph fracture morphology of steel cooled to 620 °C is shown in Fig. 8(a). The zone 1 in the macrograph stands for radical zone. The fractograph of radical zone was predominantly cleavage fracture with some obvious cracks indicated by red arrows, as shown in Fig. 8(b). Fig. 8(c) shows the macrograph fracture morphology of steel cooled to 580 °C. The zone 1 is radical zone and the zone 2 is shear lip zone. The fractograph of radical zone was predominantly quasi-cleavage facets with some small deep dimples (marked by blue arrows) and small microcracks (marked by red arrows), as shown in Fig. 8(d). Fig. 8(e) shows the fractograph of shear lip zone, which mainly consisted of small deep dimples with small precipitates (marked by pink circles) in it. In steel cooled to 620 °C,
measurement of the carbide size. The size of the TiC in Fig. 6(a) was determined to be ~4.41 nm along the direction of the Moiré fringe (length) and ~4.32 nm perpendicular to the direction of the Moiré fringe (thickness). The aspect ratio (length/thickness) is close to 1. Thus, it can be concluded that the morphology of the carbide was close to spherical. The average value of length and thickness of TiC was ~4.36 nm. Futhermore, Fig. 6(a) also reveals an excellent coherent relationship between the TiC phase and the matix. The nanoscale coherent interfaces between TiC and the matrix have a low interaface energy and good microsturcture stablity. Fig. 6(b) shows the representive HRTEM image of the Fe3C, the lattice parameter can be calculated and the results match well with the SAED pattern. Fig. 7 shows the size distribution histograms of TiC and cementite precipitated in steel cooled to 620 and 580 °C. With each speicmen at least 100 precipitates was measured. The averge sizes of TiC and cementite was 5.4 and 32.8 nm in steel cooled to 620 °C, while for the steel cooled to 580 °C, the averge size of TiC and cementite was 3.8 and 22.4 nm. 3.3. Mechanical properties Table 2 summarizes the mechanical properties of the experimental steels after hot rolling with different UFC finish temperatures, including yield strength, tensile strength, total elongation and Charpy impact energy. Each datum is obtained from three specimens. The tensile and Charpy impact test results indicate that the mechanical property reaches optimal with excellent combination of strength and toughness
Fig. 5. TEM images of steel cooled to 540 °C: (a) the cementites located on the boundaries of bainite laths and (b) the TiC precipitate formed in bainite lath; the insets in (a) and (b) are the SEAD patterns of cementite and TiC.
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Fig. 6. Representative HRTEM images of (a) TiC and (b) cementite particles.
Fig. 7. Size distribution histograms of TiC and cementite precipitated in steel cooled to (a) 580 °C and (b) 620 °C.
3.4. Strengthening mechanism
Table 2 Mechanical properties with standard deviation of the experimental steels with different UFC finish temperatures. Steel
1 2 3
Yield strength, MPa
615 (2.8) 650 (4.3) 590 (2.1)
Tensile strength, MPa
735 (5.7) 750 (2.8) 730 (6.1)
Elongation, %
22.6 (1.3) 17.4 (1.1) 10.9 (0.9)
Chemical phase analysis and SAXS analysis were carried out on steels cooled to 620 °C and 580 °C. Table 3 summarizes the structural parameters of the extracted precipitates by electrolysis using X-ray diffraction. It can be concluded from the XRD results that the precipitates were mainly composed of M3(C, N), Ti2CS, TiC and Ti(C, N). The weight fraction of MC and M3C are presented in Table 4. The content of M3C cementite-based precipitates dominated compared with that of MC. The weight fraction of M3C and MC type precipitates is similar in steel cooled 620 °C and 580 °C. The volume fraction and the corresponding strengthening increment of M3C and MC type precipitates in different diameter ranges below ~ 36 nm are shown in Table 5. It can be noted that the volume fraction of M3C and MC carbides with size less than ~36 nm in steel cooled to 620 °C is smaller than that in steel cooled to 580 °C, which will result in difference in strengthening contribution. The main objective of NG-TMCP is to enhance the strength of steel by precipitation hardening. It is therefore essential to estimate the contribution of the precipitates toward effective strengthening. There are two well-known mechanisms for precipitation strengthening, shearing mechanism and bypass mechanism [1]. The equations for the two contribution are as followings [23].
Impact toughness (V-notch, J) −20 °C
−40 °C
62.4 (7.2) 93.4 (8.6) 26.3 (3.8)
43.5 (3.6) 65.7 (6.7) 18.4 (2.4)
seldom shear lip zone can be observed, indicative of poor lowtemperature toughness compared with steel cooled to 580 °C. The main reason for the low impact energy in steel cooled to 620 °C is the existence of relatively large size cementites compared with that in steel cooled to 580 °C. Another factor that affects low temperature toughness should be related to microstructure. Fig. 9(a) and (b) shows the quality maps of steels cooled to 620 and 580 °C, respectively. The red lines represent the small angle grain boundaries with misorientation of 2–15°, while the black lines correspond to high angle grain boundaries with misoientation of 15° and larger than 15°. The high misorientation grain/packet boundaries can effectively deflect or even arrest the propagation of cleavage microcracks, whereas the low misorientation boundaries have less ability to deflect the crack [18,19]. It is obvious to notice that the percentage of high misoriention boundary in steel cooled to 580 °C is higher compared that in steel cooled to 620 °C, as shown in Fig. 9(c) and (d). These high grain boundaries in steel cooled to 580 °C contribute more resistance to cleavage microcrack [20–22].
σbypass = 10.8
f d
σshearing = M τP = 273
ln(1630d )
γ 3/2 2 × 1.1 × 2 × d1/2f 1/2 b 2AG
(1)
(2)
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Fig. 8. The impact fractographs: (a) macrograph and (b) the fractograph of zone 1 (radical zone) in steel cooled to 620 °C; (c) macrograph and the fractographs of (d) zone 1 (radical zone) and (e) zone 2 (shear lip zone) in steel cooled to 620 °C.
Fig. 9. EBSD image quality maps with grain boundary misorientation distribution of steels with UFC finish cooling temperature of (a) 620 °C and (b) 580 °C; grain boundary distribution histograms of steels with UFC finish cooling temperature of (c) 620 °C and (d) 580 °C.
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Table 3 Structural parameters of the precipitates.
Table 6 Comparison of calculated yield strength values with experimental values in steels with UFC finish temperatures of 620 and 580 °C.
Phase structure
Lattice constant, nm
M3(C,N)
a0=0.4523–0.4530, b0=0.5088–0.5080, c0=0.6743–0.6772 a0=0.3210–0.3240， c0=01.1203–1.1308， c/a=3.49 a0=0.431–0.433 a0=0.425–0.427
Ti2CS TiC Ti(C,N)
Crystal system No.
D, μm
Orthorhombic 1 2
Hexagonal
7.5 6.9
Calculated Yield Strength Values, MPa σGS
σSS
σSP
σy
*σSP
*σy
Actual measured σy, MPa
218.9 228.3
77.7 78.6
279.4 306.9
576.0 613.8
314.4 351.6
611.0 658.5
615 650
Face centered cubic Face centered cubic
Table 4 The mass fraction and chemical structural formula in steel with UFC finish temperature of 620 and 580 °C. No.
1 2
MC phase
M3C phase
Mass fraction (%)
Structural formula
Mass fraction (%)
Structural formula
0.0313 0.0338
Ti(C0.609N0.391) Ti(C0.672N0.328)
1.4113 1.5553
(Fe0.9856Mn0.0144)3C (Fe0.984Mn0.016)3C
Where, τP represents the shear stress caused by dislocations cutting the particles in MPa; b is dislocation Burger's vector and equal to 0.248 nm; G is the shear elasticity modulus and equal to 80650 MPa; γ is the interface energy between precipitates and matrix and equal to 0.5–1 J/ m2; d represents the second-phase particle diameter in μm; f represents the volume fraction of the precipitates; From formula (1) and (2), we can notice that the strengthening increment of shearing mechanism increases with the increasing precipitate size, while the strengthening increment of the bypass mechanism decreases with the increase of particle size. The critical transformation size dc can be calculated by the following equation.
Gb 2 d ln( c ) dc = 0.209 Kγ 2b
Fig. 10. The magnetic scattering of steels cooled to 620 and 580 °C by SANS. (Measured data are presented in solid square or circle, fitted data are presented in solid line).
σy = σSG + σSS + σSP = 600D−1/2 + {46[C ] + 37[Mn] + 83[Si] + 59[Al] + 2918[N ] + 80.5[Ti]} + σSP
(4)
where σy , σSG , σSS and σSP represent the yield strength, grain refinement strengthening, solid solution strengthening, and precipitation strengthening in MPa, respectively. The components of the yield strength and the total yield strength are presented in Table 6. The result shows that estimated yield strength for steels cooled to 620 and 580 °C are 576.0 MPa and 613.8 MPa, lower than the actual yield strength. Exploring the reason, it should be attributed to the limitation of electrolytic separation that some fine precipitates lost during the electrolyzed powders filter process. So, the SANS was chosen to confirm the volume fraction of tiny precipitates with bulk samples. Fig. 10 shows the magnetic scattering curves of steels cooled to 620 °C and 580 °C. The magnetic scattering intensity was extracted from the intensity collected when the scattering vector is perpendicular (α=90°) and parallel (α=0°) to the applied magnetic field, according to
(3)
The meaning of symbols in Eq. (3) is same as that mentioned above. The critical transformation size of the precipitate depends on the interface energy between the precipitates and the matrix. The dc of TiC and Fe3C were estimated to be ~1.5–6 nm and ~4.7–10 nm calculated by Eq. (3). The precipitation strengthening effect of TiC in all size ranges and Fe3C with size large than 10 nm were calculated based on the bypass mechanism. For Fe3C with size less than 10 nm, the shearing mechanism was applied [12]. The calculated results of precipitation strengthening increment in different size range are listed in Table 6. For low-carbon steel, the yield strength equals to the sum of solid strengthening, grain refinement strengthening, and precipitation strengthening, and is given by the following equation [12].
Imagnetic = I (α = 90°) − I (α = 0°)
(5)
Table 5 Contribution of Fe3C and TiC to the yield strength of steels with UFC finish temperatures of 620 and 580 °C. No.
Diameter Range, nm
Fe3C
TiC
Total increment, Mpa
Volume Fraction, Pct
Yield Strength Increment, MPa
Volume Fraction, Pct
Yield strength increment, MPa
1
1–5 5–10 10–18 18–36 ∑
0.0153 0. 0038 0.0020 0.0025 0.0236
60.5 16.5 7.6 5.2 89.7
0. 0364 0 0. 0873 0. 3550 0.4786
78.6 0.0 50.1 61.0 189.7
279.4
2
1–5 5–10 10–18 18–36 ∑
0.0060 0.0026 0.0011 0.0032 0.0130
38.0 13.6 5.6 5.9 63.1
0.1172 0.0035 0.0019 0.2226 0.3452
153.9 34.2 7.4 48.3 243.8
306.9
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3. The contribution of precipitates to yield strength was greater than ~350 MPa steel with UFC finish cooling temperature, which clearly indicates that precipitation strengthening primarily contributed to the total yield strength. A small amount of Ti (0.08 wt%) addition can improve the traditional Q345 steel to obtain 650 MPa in yield strength.
Table 7 The fine precipitate size and volume fraction in steels with UFC finish temperatures of 620 and 580 °C measured by SANS. No.
1 2
SANS Rs (nm)
Gs
f
2.4(0.2) 1.8(0.1)
0.09(0.02) 0.17(0.03)
0.055% 0.130%
Acknowledgements The research was supported financially by National Science Foundation of China (Grant No. 51234002, 51504064 and 51474064) and National Key Research and Development Program 2016YFB0300601. The China Scholarship Council is also acknowledged sincerely for supporting us to complete this work. RDKM gratefully acknowledges support from University of Texas at El Paso, USA.
Three different slopes can be observed for both samples. The low-q data decays with q−4, which is induced by grain boundaries and large precipitates. The slope of the median- and high-q data is −2.7 and −1.8 in logarithmic coordinate, respectively. It means that the precipitates have two characteristic sizes. We fit the data with a power law and Guinier laws due to the limited fitting range,
I (q ) = Gs exp(−
q 2Rg2 q 2Rs2 ) + G exp(− ) + Aq−4 3 3
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where, Rs and Rg are the size-weighted Guinier radiuses of precipitates with small and medium size, GS and G are the scale factors. The last term in Eq. (6) accounts for the scattering from large precipitates and grain boundaries. The fitting parameters are listed in Table 7. The average radius and the volume fraction of the tiny precipitates (1–5 nm) in steel cooled to 620 °C are 2.4 nm and 0.055%. For steel cooled to 580 °C, they are 1.8 nm and 0.130%, respectively. The precipitation strengthening increment in precipitates (1–5 nm) was estimated to be 174.1 and 198.6 MPa in steels cooled to 620 and 580 °C, respectively. The total precipitation strengthening (*σSP) for the two steels were 314.4 and 351.6 MPa, respectively. The calculated yield strength (*σy) matches well with the measured yield strength, as shown in Table 6. It can be concluded that the contribution of precipitates to yield strength can be greater than ~350 MPa in steel cooled to 580 °C, which clearly indicates that precipitation strengthening primarily contributed to the total yield strength. 4. Conclusion 1. Besides nanoscale TiC, cementite precipitates of size less than ~36 nm were also observed in Ti-microalloyed steel. The volume fraction of Fe3C was significantly higher than TiC for similar size range. 2. When the steel was ultrafast cooled to 580 °C, combination of high yield strength of 650 MPa, good elongation of 17%, and excellent toughness of 90 J at −40 °C was obtained through nanoscale cementites and microalloyed TiC carbides precipitation strengthening.
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