Regular Article
Suppression of Coarsening of Upper Bainite at Prior Austenite Grain Boundaries by Intra-granular Bainite Nucleated on Inclusions in Steel Welds
2024 Volume 64 Issue 1 Pages 126-133
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2024 Volume 64 Issue 1 Pages 126-133
The microstructures of intra-granular bainite nucleated on Ti-oxide particles (IGB) and grain boundary upper bainite (GBB) in Ti-deoxidized steel (Fe-0.1mass%C-1.5mass%Mn-2mass%Ni-1mass%Cu) were investigated by EBSD analysis and 3-dimentional observation. The steel was austenitized at 1673 K for 23 s, held at 803 K for 7 s-20 ks and then quenched to room temperature. 803 K is just below the bainite transformation start temperature. IGB and GBB were observed at 5% bainite transformation. Despite the formation of IGB, part of GBB grew to a size of 100 µm at 17% bainite transformation, and coarsened to 130–200 µm at 85% bainite transformation. Mechanism of the GBB coarsening is discussed in terms of differences in (1) the microstructures and (2) the nucleation site of IGB and GBB. A single packet with many blocks was observed in GBB nucleated at whole surface of austenite grain boundaries with a size of 400–500 µm, while multiple packets with two blocks were observed in IGB nucleated on Ti-oxide particles with a size of 1–2 µm. IGB suppressed the growth of GBB by impingement. However, as the GBB was much larger than IGB, part of the GBB was not impinged by the IGB and continued to grow. GBB grew until all blocks of GBB were impinged by IGB and/or transformation was finished, resulted in GBB coarsening.
In recent years, welded structures have become larger and lighter; as a result, steel plates that support them are required to have high strength. Furthermore, steel plates must have high toughness to ensure the reliability and safety of the welded structures. One notable challenge in the application of high-strength steels is ensuring the toughness of the heat-affected zone (HAZ).
Typically, the steel plates are exposed to a minimum temperature of 1350°C during welding, which coarsens the microstructure in the HAZ, and consequently deteriorates the toughness of the HAZ. An effective approach for improving the HAZ toughness is to refine the crystal grains.1,2,3) One method proposed to accomplish this is the utilization of an intra-granular transformation microstructure,3) which nucleates from inclusions in austenite (γ) grains. The formation of intra-granular ferrite (IGF) from titanium (Ti) oxide has often been reported in the HAZs of low-strength steels (up to TS690 MPa). Moreover, the formation mechanism of IGF4,5,6) and the toughness improvement due to IGF7,8) have been clarified.
By contrast, in the HAZ of high-strength steel (over TS700 MPa), the main microstructure changes from ferrite to upper bainite. Although grain boundary ferrite (allotriomorph) is formed by a diffusion mechanism, grain boundary upper bainite (GBB) is formed by a shear mechanism and grows rapidly; therefore, the crystal grains coarsen in the HAZ of high-strength steels.9,10,11) Furthermore, GBB has been characterized as a substructure with a small misorientation angle and does not interfere with the propagation of cleavage cracks, so cracks propagate rapidly and the toughness deteriorates.12) Suppressing the coarsening of GBB and refining the crystal grains are effective strategies for improving the HAZ toughness of high-strength steels.
Intra-granular transformation, similar to its application in low-strength steel, can be considered a means of suppressing the coarsening of GBB within the HAZs of high-strength steels. However, few studies have focused on the intra-granular transformation of HAZs in high-strength steels. Shinohara et al.13) reported that intra-granular bainite (IGB) nucleated from Ti oxides in the HAZs of high-strength steels and that IGB was lath-like and possessed a similar morphology to GBB. Additionally, Shinohara et al.14) reported that IGB and GBB shared a common {110}α plane and that the {110}α planes of IGB and GBB were almost parallel. However, how the microstructures and transformation temperatures differ between IGB and GBB are unknown.
Shigesato et al.15) confirmed the presence of a Mn-depleted zone surrounding the Ti oxides formed by IGB. Moreover, they reported that IGB formed owing to an increase in the bainite transformation temperature around the inclusion. Similar studies have been conducted for IGF,4,5,6) which has the same formation mechanism as IGB.
Studies13,16) have indicated that the formation of IGF and IGB in the HAZ or weld metal of high-strength steel improves toughness owing to microstructural refinement. However, Nako et al.17) and Lan et al.18) confirmed the formation of relatively coarse GBB despite the formation of IGF and IGB, and suppressing the production of GBB itself was not possible. To refine the grain size of the HAZ in high-strength steels using IGB, it is necessary to clarify the differences in the microstructures, transformation temperatures, and competitive states of IGB and GBB.
In this study, we held Ti-deoxidized steel equivalent to the TS780 MPa class at 530°C, near the bainite transformation starting temperature, and then quenched it to obtain specimens. We then observed their IGB and GBB microstructures in the early transformation stages using scanning electron microscopy (SEM), determined their crystal orientation using electron backscatter diffraction (EBSD), and observed their three-dimensional (3D) structure using optical microscopy and serial sectioning.
Table 1 presents the chemical compositions of the tested steels. The basic composition was Fe–0.1C–1.5Mn–2Ni–1Cu (mass%), which is characteristic of the TS780 MPa class. Ti-deoxidized steel was prepared via vacuum melting. Subsequently, it was heated at 1150°C for 3600 s and then hot-rolled to obtain 15 mm thick plates. A round-bar specimen with dimensions of φ 3 mm × 10 mm (with a hole of φ 2 mm × 4 mm to attach the thermocouple) was extracted from the rolled material.
| Chemical composition [mass%] | Temperature [°C] | |||||||||||
|---|---|---|---|---|---|---|---|---|---|---|---|---|
| C | Si | Mn | P | S | Ni | Cu | Al | Ti | N | O | Bs | Ms |
| 0.1 | 0.1 | 1.5 | 0.01 | 0.002 | 2.0 | 1.0 | <0.002 | 0.011 | 0.004 | 0.001 | 540 | 430 |
We conducted a preliminary continuous-cooling test in which the specimen was heated at 1400°C for 23 s and then cooled from 1000°C to 20°C at a rate of 0.1–100°C/s. Based on the microstructure and hardness at each cooling rate, the bainite transformation start temperature (Bs point) and martensite transformation start temperature (Ms point) were determined to be 540 and 430°C, respectively.
2.2. Isothermal-holding TestThe specimen was heated at 1400°C for 23 s to simulate the prior γ grain size (400–500 μm) of welds corresponding to a heat input of 50 kJ/mm. Afterwards, in order to simulate the HAZ microstructure of the high-strength steel, grain boundary ferrite formation was suppressed. To create a single-phase bainite structure, the specimen was quenched by helium (He) gas until 530°C; after holding for 7 s, 21 s, and 20 ks at 530°C, just below the Bs point, the specimen was quenched with He gas until 20°C (Fig. 1). We utilized this heat-treatment method to obtain bainite microstructures at the early stage of transformation by holding for 7 and 21 s at 530°C and at the end stage of transformation by holding for 20 ks at 530°C. Subsequent quenching resulted in the formation of martensite from the untransformed austenite. The cooling rate during quenching was set to 100°C/s, which was confirmed in the preliminary test to induce a martensite phase. Induction heat treatment was conducted using a Formastor tester (Fuji Electric Koki Co., Ltd.), and the changes in dilatation with respect to temperature were measured. The transformation rate was calculated by dividing the dilatation at each holding time by the total dilatation.
The cross-sections of the isothermally held and quenched specimens were mirror-polished and etched with a 3% nital solution. Subsequently, the microstructures were examined using a field-emission scanning electron microscope (FE-SEM; JEOL Ltd., JSM-6500F) at an accelerating voltage of 15 kV and magnification of 250×–4000×. The obtained SEM images were used to measure the GBB area, and the GBB size was calculated using the circle-equivalent diameter.
Additionally, the elemental composition of the inclusions, which were the nucleation sites of IGB, was analyzed using energy-dispersive X-ray spectroscopy (EDS) with an instrument attached to the FE-SEM.
2.3. Crystal Orientation AnalysisAfter mirror-polishing the cross section of the specimen, which was held at 530°C for 21 s and then quenched, electrolytic polishing was performed using a chromate-phosphate solution as the electrolyte. The crystal orientations of IGB and GBB were measured using EBSD at a magnification of 2500×–3000× and a step width of 0.1 μm.
Table 2 presents 24 variants of the Kurdjumov–Sachs (K–S) orientation relationship.19) The close-packed (CP) groups are denoted as CP1–CP4, and the groups with the same Bain correspondence are denoted as B1–B3. Variants of the same Bain correspondence group were generated around the {001}γ of prior γ grains, as indicated by the (001) pole figure illustrating the K–S relationship (Fig. 220)). In this study, the boundary with a misorientation of 5° was defined as the block boundary.21) The block widths of IGB and GBB were measured as the length vertical to the growth direction. Additionally, a rotation operation was performed in the (001)γ plane to match the 24 variants of the K–S relationship and pole figures obtained from the EBSD analysis.
| Variant | Plane parallel | Direction parallel | Angle from V1 | Group |
|---|---|---|---|---|
| V1 | (111)γ// (011)α’ CP1 | [-1 0 1]γ // [-1 -1 1]α’ | – | B1 |
| V2 | [-1 0 1]γ// [-1 1 -1]α’ | 60.0 | B2 | |
| V3 | [0 1 -1]γ// [-1 -1 1]α’ | 60.0 | B3 | |
| V4 | [0 1 -1]γ// [-1 1 -1]α’ | 10.5 | B1 | |
| V5 | [1 -1 0]γ// [-1 -1 1]α’ | 60.0 | B2 | |
| V6 | [1 -1 0]γ// [-1 1 -1]α’ | 49.5 | B3 | |
| V7 | (1-11)γ// (011)α’ CP2 | [1 0 -1]γ // [-1 -1 1]α’ | 49.5 | B2 |
| V8 | [1 0 -1]γ// [-1 1 -1]α’ | 10.5 | B1 | |
| V9 | [-1 1 0]γ// [-1 -1 1]α’ | 50.5 | B3 | |
| V10 | [-1 1 0]γ// [-1 1 -1]α’ | 50.5 | B2 | |
| V11 | [0 1 1]γ// [-1 -1 1]α’ | 14.9 | B1 | |
| V12 | [0 1 1]γ// [-1 1 -1]α’ | 57.2 | B3 | |
| V13 | (-111)γ// (011)α’ CP3 | [0 -1 1]γ // [-1 -1 1]α’ | 14.9 | B1 |
| V14 | [0 -1 1]γ// [-1 1 -1]α’ | 50.5 | B3 | |
| V15 | [-1 0 -1]γ// [-1 -1 1]α’ | 57.2 | B2 | |
| V16 | [-1 0 -1]γ// [-1 1 -1]α’ | 20.6 | B1 | |
| V17 | [1 1 0]γ// [-1 -1 1]α’ | 51.7 | B3 | |
| V18 | [1 1 0]γ// [-1 1 -1]α’ | 47.1 | B2 | |
| V19 | (11-1)γ// (011)α’ CP4 | [-1 1 0]γ // [-1 -1 1]α’ | 50.5 | B3 |
| V20 | [-1 1 0]γ// [-1 1 -1]α’ | 57.2 | B2 | |
| V21 | [0 -1 -1]γ// [-1 -1 1]α’ | 20.6 | B1 | |
| V22 | [0 -1 -1]γ// [-1 1 -1]α’ | 47.1 | B3 | |
| V23 | [1 0 1]γ// [-1 -1 1]α’ | 57.2 | B2 | |
| V24 | [1 0 1]γ// [-1 1 -1]α’ | 21.1 | B1 |
Observation, imaging, polishing, and etching of the surface of the specimen, which was held at 530°C for 21 s and then quenched, were automatically performed using an optical microscope and the serial sectioning method (Genus_3D, Nakayama Electric Co., Ltd.). The obtained image data was used to reconstruct 3D images. The automatic polishing depth was approximately 0.5 μm per round, and Avizo software for industrial inspection (Thermo Fisher Scientific K. K.) was used for 3D construction.
Figure 3 shows the changes in dilatation during the isothermal-holding tests. The transformation rates at 530°C for holding times of 7 s, 21 s, and 20 ks were 5%, 17%, and 85%, respectively.
Figure 4 shows the nital-etched microstructure of the specimen (transformation rate: 85%) quenched after holding at 530°C for 20 ks. As explained later in the manuscript, the IGB formation was confirmed, and the microstructure of the prior γ grains consisted of fine IGB. Additionally, a relatively coarse GBB with a circle-equivalent diameter of 168 μm formed from the prior γ grain boundary. The circle-equivalent diameter of the GBB observed in six different fields ranged from 130 to 200 μm. However, if the entire microstructure consisted of GBB and no IGB, the GBB size was expected to be approximately 400 μm, which was equivalent to the prior γ grain size in this study. By contrast, for Ti-deoxidized steel containing IGB, the GBB size was relatively small, and the microstructure of the prior γ grain was fine. Therefore, the utilization of IGB clearly improves the toughness. These results agree with those of previous studies.13,16)
Figure 5 shows the nital-etched microstructure of a specimen (transformation rate: 5%) quenched after holding at 530°C for 7 s. The microstructure consists of bainite, which formed during heat treatment, and martensite, which austenite transformed into after quenching. At the start of the transformation, an intra-granular transformation microstructure nucleated on the inclusion (Fig. 5(a)). The intra-granular transformation occurred through a shearing mechanism because the microstructure was lath-like and the transformation temperature was low in steel with high hardenability. Similar microstructures were observed in a previous study,13) and this intra-granular transformation microstructure was identified as IGB. In addition to IGB, GBB formed (Fig. 5(b)) at the start of the transformation. In this study, the bainite formed during the 530°C holding consisted of IGB and GBB. According to the results of the EDS elemental analysis of five inclusions with confirmed IGB formation, the nucleation sites of IGB were inclusions containing Ti oxides. These results agree with those of previous studies.13,16)
Figure 6 shows the nital-etched microstructure of the specimen (transformation rate: 17%) quenched after holding at 530°C for 21 s, which was done to observe the microstructure at the early stage of transformation. At this stage, GBB had already grown to a circle-equivalent diameter of 100 μm and length of 90 μm. However, the growth of a portion of GBB was delayed, and the length of this portion was 60 μm (Fig. 6(a)). The high-magnification SEM images (Fig. 6(b)) revealed multiple IGB formations in regions where the growth of GBB was delayed.
Figures 7 and 8 show the SEM and crystal orientation analysis results for two IGB groups (IGB① and IGB②) in a specimen that was held at 530°C for 21 s and then quenched. Areas with a small misorientation (<5°) (corresponding to one block) are highlighted and colored in the IQ maps shown in (b) and pole figures shown in (c). Furthermore, the pole figures (Figs. 7(c) and 8(c)) are rotated for comparison with the 24 variants of the K–S relationship (Figs. 7(d) and 8(d)). In this context, the variant of IGB highlighted in green was defined as V1, and the other variants were analyzed.
Blocks formed in multiple directions from one inclusion in both IGB① and IGB②. The block width of IGB was 1–3 μm. In IGB①, blocks of the same CP group (V3 or V4 relative to V1) nucleated in the same direction and formed a packet. In addition, blocks of the same Bain correspondence group (V8 relative to V1) with small misorientations were observed in the same direction. Blocks of different CP groups (V23 relative to V1) formed in another direction. It has been reported that V1/V4 and V1/V8 pairs with equally small misorientations are frequently observed22) and that the stress and strain associated with the transformation are relaxed23) in martensite, with similar behavior observed in IGB. Similarly, in IGB②, as with IGB①, blocks of the same CP group (V2 relative to V1) and blocks of the same Bain correspondence group (V8 relative to V1 and V7 relative to V2) with small misorientations formed in the same direction. In this study, IGB was observed in five fields, all of which possessed similar crystallographic features.
3.2.2. GBB Crystal OrientationFigure 9 shows the SEM and crystal orientation analysis results for GBB in a specimen that was held at 530°C for 21 s and then quenched. As mentioned in Section 3.2.1, an area with a small misorientation (<5°) is highlighted and colored in the IQ map shown in (b) and pole figure shown in (c), and the pole figure is rotated. In GBB, many blocks formed discontinuously from the prior γ grain boundary and were all the same variant. The block width of GBB was 1–3 μm, which was identical to that of IGB. GBB was observed in five fields, all of which possessed similar crystallographic features.
Figure 10 shows the SEM and crystal orientation analysis results for IGB and GBB in a specimen that was held at 530°C for 21 s and then quenched within the same field, as shown in Fig. 6. As mentioned in Section 3.2.1, areas with a small misorientation (<5°) are highlighted and colored in the IQ map shown in (a) and pole figure shown in (b), and the pole figure is rotated. In this context, the GBB variant highlighted in red was defined as V1, and the IGB variants were analyzed.
As shown in Fig. 9, GBB was composed of blocks of the same variant and consisted of a single packet. The six IGB observed in Fig. 10 belonged to a different CP group than GBB. The measurement results for the crystal orientation relationship of IGB and GBB in six fields indicated that among the 41 IGB variants, 6 belonged to the same CP group as GBB, and 35 belonged to a different CP group than GBB. The {110}α planes of IGB and GBB were almost parallel and belonged to the same CP group in a previous study,14) but many of the IGB variants belonged to a different CP group than GBB in the present study. Some IGB, which belonged to a different CP groups than GBB, existed in the growth direction of the GBB.
The same V1 variant microstructure exhibited by GBB existed on the left side of IGB in V15, V17, and V18 variants in Fig. 10(a). GBB growth was delayed in Fig. 6; however, the crystal orientation analysis indicated that GBB grew through IGB.
3.3. 3D Structural ObservationAs mentioned in Section 3.2.3, it was possible that GBB grew through IGB; therefore, to clarify the 3D positional relationship between IGB and GBB, a 3D image was reconstructed using optical microscopy and serial sectioning. Figure 11 shows optical micrographs of IGB and GBB before reconstruction of the 3D image in a specimen held at 530°C for 21 s and then quenched. Similar to Fig. 6, IGB③ and IGB④ existed in the growth direction of GBB. Figure 12 presents 3D images of IGB and GBB. Only the inclusion, IGB, and GBB are colored without the martensite that transformed during quenching. IGB③ and IGB④ were nucleated on one inclusion in different directions. As shown in Fig. 11, GBB growth stopped at the point of collision with IGB. IGB suppressed the growth of GBB via hard impingement; however, GBB three-dimensionally formed in the same direction along the entire plane of the prior γ grain boundary, and it continued to grow in the part that did not collide with IGB.
Table 3 presents the differences in the microstructures of IGB and GBB, including schematics. In IGB, two blocks with widths of 1–3 μm are nucleated on one inclusion in the same direction, forming a packet. In addition, blocks of the same Bain correspondence group with small misorientations may form. IGB forms in multiple directions, resulting in multiple packets. By contrast, GBB forms as many blocks with widths of 1–3 μm from the prior γ grain boundary; these blocks are of the same variant and constitute a single packet.
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According to Kaneshita et al.,24) the variant selection rules for GBB nucleation are constrained by the following four conditions: (1) the bainite has a near K–S relationship with γ grains on both sides; the bainite (2) growth direction and (3) habit plane are parallel to the γ grain boundaries; and (4) plastic accommodation of the bainite transformation strain, where constraint (1) is the strongest, particularly with a low C content and high transformation temperature. Figures 6 and 9 show that GBB grows on both sides of the γ grain, where GBB is composed of blocks of the same variant because it has a near K–S relationship with the γ grains on both sides (constraint (1)). By contrast, IGB is not constrained by (1), which is the strongest rule among the four. Additionally, since the nucleation site is curved at the interface between the inclusion and the matrix, IGB can select the variant in which the bainite (2) growth direction and (3) habit plane are parallel to the interface. Thus, IGB was composed of multiple variants. The different number of blocks constituting the IGB and GBB groups was attributed to the size of each nucleation site. In IGB, many blocks could not be nucleated from inclusions, because the inclusions of the nucleation sites were only a few micrometers—nearly identical to the width of the IGB blocks. By contrast, the 3D structure of GBB contained many adjacent blocks along a wide γ grain boundary. The nucleation site of GBB comprised the entire plane of the prior γ grain boundary, whose size was 400–500 μm, and many blocks formed because the nucleation site was considerably larger than the GBB block.
These results indicate that the IGB and GBB microstructures differ owing to their different nucleation behaviors; different constraints of variant selection rules in bainite nucleation, as well as the different number of adjacent blocks arising from the differences in their nucleation sites.
4.2. Transformation Temperature and Competitive States of IGB and GBBFigure 13 shows a schematic of the bainite transformation process. At the start of the transformation, both IGB and GBB form and compete with each other (Fig. 13①). At the early stages of transformation, GBB forms as many blocks of the same variant from the γ grain boundary, resulting in coarsening to approximately 100 μm. By contrast, GBB growth is suppressed by hard impingement of IGB (Fig. 13①). However, the 3D observations revealed that GBB stopped growing in the part that collided with IGB but continued to grow in the part that did not collide. Therefore, GBB grows until all blocks collide with IGB and/or the transformation is complete, and it coarsens to a circle-equivalent diameter of 130–200 μm (Fig. 13②).
Based on these results, which are in accordance with those of previous studies, relatively coarse GBB formed despite IGB formation in Ti-deoxidized steel because GBB grew until either all blocks collided with IGB or the transformation was complete. The grain sizes of IGB and GBB differ at the early stage of transformation because GBB forms in more blocks than IGB, and part of the GBB does not collide with IGB. Therefore, to suppress the coarsening of GBB and further refine the grains of the HAZ of high-strength steel by utilizing IGB, it is necessary to increase the number of collisions between IGB and GBB during the early stage of transformation.
In addition, the driving force for IGB nucleation is important for the effective utilization of IGB. In general, the transformation starts from the γ grain boundary.3,25) However, in this study, the transformation temperatures for IGB and GBB were identical. This is because the driving force for α transformation increases owing to the formation of the Mn-depleted zone around the Ti oxides,4,5,6,15) and regarding the thermal history in this study, the transformation temperature of IGB may have increased to that of GBB. If the driving force for IGB nucleation can be further increased, the transformation temperature of IGB can be higher than that of GBB. As a result, the number of IGB formed during the early stages of transformation would increase, and enhanced suppression of GBB coarsening would be possible.
To clarify the difference between the IGB and GBB microstructures, transformation temperatures, and competitive states, we held TS780 MPa class Ti-deoxidized steel at 530°C (near the Bs point) and then quenched it. We then analyzed the characteristics of IGB and GBB at the early stage of transformation by observing their microstructures using SEM, their crystal orientation using EBSD, and their 3D structures using optical microscopy and serial sectioning. Our findings are summarized below.
(1) Regarding the differences between the IGB and GBB microstructures, IGB formed multiple packets of two blocks from the interface between the inclusion and matrix. By contrast, GBB formed a single packet of many blocks of the same variant along the entire plane of the coarse γ grain boundary. GBB consisted of blocks of the same variant because the variant selection rule was constrained owing to the near K–S relationship with γ grains on both its sides. Accordingly, IGB and GBB microstructures differ owing to their different nucleation behaviors; different constraints of variant selection rules in bainite nucleation, as well as the different number of adjacent blocks arising from the differences in their nucleation sites.
(2) Regarding the competitive states of IGB and GBB, the transformation temperatures of IGB and GBB were identical. In GBB, at the early stages of transformation, many blocks of the same variant were formed from the prior γ grain boundary, resulting in coarsening to approximately 100 μm. By contrast, GBB growth was suppressed by hard impingement of IGB. However, GBB stopped growing in the part that collided with IGB but continued to grow in the part that did not collide. Relatively coarse GBB formed despite IGB formation in Ti-deoxidized steel because GBB grew until either all blocks collided with IGB or the transformation was complete.
(3) To suppress the coarsening of GBB and refine the HAZ grains of high-strength steel by utilizing IGB, it is necessary to increase the number of collisions between IGB and GBB at the early stage of transformation and/or to increase the driving force for IGB nucleation.
