Figure 3-2 shows the optical metallography of each sample. The both microstructures consist of ferrite and pearlite and the grain size of the ferrite is relatively fine (< 10 μm). The difference in grain sizes of ferrite was mainly controlled by the degree of deformations. For Steel A with more amount of thermo-mechanical deformation, the grain size of ferrite was further to be smaller than 5 μm. For Steel B, the banding structure can be observed. The existence of the banding structure was supposed due to element segregation, especially those of manganese and sulfur.
However, in the later section, the testing of the mechanical properties showed that the banding structure did not have any significant effect on the mechanical properties. The grain size and volume fraction of each phase are listed in Table 3-2.
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Figure 3-2 Optical micrographs of UHSSBs revealing microstructures composed of ferrite and pearlite:
(a)Steel-A; (b) Steel-B.
Table 3-2 The microstructural features of UHSBs determined by TEM
Sample No. Steel-A Steel-B
Ferrite volume fraction (%) 40.3 45.2
Ferrite grain size (μm) 4.7 1.2± 8.3 0.8±
Sheet spacing of interphase precipitation (nm) 19.6±2.8 20.3 1.8±
Carbide size (nm) 9.0 2.3± 11.4 2.6±
Pearlite spacing (nm) 173.4 11.7± 145.8 8.3±
Dislocation density in ferrite (m-2) 9.3 10× 13 2.1 10× 13
b. Mechanical properties
In the beginning, Vickers microhardness test has been used for the estimation of the strength of each phase. Because the grain size of ferrite is relatively fine, the indentation would cover the grain boundaries or very close to grain boundaries. Thus, the effects of ferrite grain boundary strengthening would be included in the tested values. In Table 3-3, the value of Vickers hardness (Hv) of ferrite in Steel A is slightly higher as compared with that in Steel B. The difference is suggested to be resulted from different ferrite grain size. The smaller the ferrite grain size, the higher the value of Hv.
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On the other hand, the Hv of the pearlite is higher in Steel B than in Steel A. That could be explained by the finer inter-lamellar spacing of pearlite in Steel B. The detailed microstructures will be discussed in the later section.
Table 3-3 The mechanical properties of UHSBs
Steel-A Steel-B
Ferrite (Hv) 281.6 7.3± 273.9 3.7±
Pearlite (Hv) 351.5 8.6± 361.2 5.9±
Yield Strength (MPa) 830.75 13.7± 752.5 10.6± Tensile Strength (MPa) 1002.6 14.2± 921.7 9.2±
Elongation (%) 11 12
Tensile tests were used to reveal the elasto-plastic behaviors of UHSB. In Figure 3-3, the relative stress-strain curves show the amount of strain of the yield plateau can rise to be about 1.4% as the stress (785 MPa) is reached. It is suggested that the proportion of ferrite was sufficient to provide soft phase to be sustaining deformed under the high stress, leading to apparent yield points and yield plateaus in stress-strain curves.
Figure 3-3 (a) Stress-Strain Curves of Steel-A and Steel-B, demonstrating apparent yield point and yield plateaus, and (b) local magnification at yield point.
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The microstructures of UHSB were similar to conventional steel bars; both consisted of ferrite and pearlite. But the results of tensile tests and macroscopic mechanical properties (presented in Table 3-3) demonstrate that the yield strengths and microhardness values of each phase are much higher than those of conventional steel bars. For conventional steel bars, the strengths of ferrite and pearlite are about 120 Hv and 280 Hv, respectively. In order to understand the increment of strengths, Transmission electron microscopy has been utilized to examine the microstructures of UHSBs in detail.
c. Transmission Electron Microscopy (TEM)
The detailed microstructure was analyzed by TEM. Figure 3-4 clearly displays that the carbides interphase precipitated in a ferritic matrix with regular spacing. Such a carbide array has been widely studied and applied to manufacture high strength steel plates for automobiles [76]. For Steel-A, the carbide size and sheet spacing are about 9 nm and 20 nm, respectively; for Steel-B, the two parameters were about 11.5 nm and 21.4 nm, respectively. It indicates that the degree of deformation before austenite-to-ferrite transformation did not affect greatly on the interphase-precipitated carbides. The carbides distributed in this manner are effective to limit dislocation motion. Figure 3-4(a) clearly gives an example showing how the interphase-precipitated carbides interact with the dislocation.
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Figure 3-4 The TEM micrographs showing typical inter-phase precipitated carbides in ferrite; specimens were deformed to different diameters and subsequently continuous cooling to the room temperature: (a) Steel-A; (b) Steel-B. Fig. 2(a) clearly showing that dislocations were pinned by these carbides.
The structure and chemistry of the interphase-precipitated carbide were analyzed by electron diffraction patterns and EDX analysis. The bright-field and dark-field image with the corresponding diffraction patterns are shown in Figure 3-5. The diffraction patterns reveal that the carbide had an FCC structure and obey one variant of Baker-Nutting orientation relationship with the ferritic matrix as follows:
(
0 0 1) (
VC || 0 0 1)
α[
1 1 0] [
VC || 1 0 0]
α, which is consistent with previous studies on vanadium-containing steels [6, 7, 17, 18].
EDX analysis shows that these carbides contain a significant level of vanadium and niobium, as indicated in Figure 3-6.
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Figure 3-5 Bright-field and dark-field images and corresponding diffraction patterns of carbides in Steel-A. (a) Bright-field image; (b) Dark-field image illuminated by 002 carbide reflection; (c) Corresponding diffraction; (d) Identified orientation relationship of carbides and ferritic matrix.
Figure 3-6 EDX analysis on nano-sized carbides in Steel-A demonstrating that they were composed of niobium and vanadium.
In addition, the lamellar spacings of the pearlite in Steels-A and -B have been resolved by TEM. Figure 3-7 shows that the lamellar spacings of pearlite are very fine,
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and no carbides were observed into the pearlitic ferrite. This result is not similar to that reported in other studies [6, 47, 52], which asserted that carbides were found in pearlitic ferrite. The occurrence of carbides in pearlitic ferrite has not been systematically investigated yet. It was presumed that the rate of transformation should be slow enough to allow carbides to have time to precipitate in pearlitic ferrite. For the present as-received materials, the prior austenite grain size was relatively small due to the addition of Nb and great amounts of defects created during hot-rolling process, so the rate of transformation of these UHSB was expected to be relatively higher. Thus, no carbides precipitated in pearlitic ferrite can be observed in the Steels-A and -B.
Figure 3-7 The fine inter-lamellar spacing of pearlite in UHSSBs: (a) Steel-A; (b) Steel-B. It shows that the spacing is about 100 nm.
The mean lamellar spacing has been measured via by mean truce spacing , λα, which is related with mean intercept spacing L. L and λα are related by a factor of 0.5 [54]:
λ =α 0.5L (3-1)
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The results of estimated lamellar spacings have been listed in Table 3-2. The lamellar spacing of Steel-B is finer than that of Steel-A. As a consequence, the measured Hv of Steel-B is higher than that of Steel-A. The difference of pearlite spacing is presumed to be resulted from the degree of deformation during the thermo-mechanical process. The lamellar spacing of pearlite depends only on transformation temperature [77]. It is proposed that Steel-A is expected to have a higher transformation rate than Steel-B because it experienced a more severe deformation. During the continuous cooling, the pearlite reaction in Steel-A would start at a higher temperature than in Steel-B, resulting in a larger pearlite spacing.