Evolution of inclusions in DH36 grade ship plate steel during high heat input welding

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Scientific Reports volume 14, Article number: 18921 (2024 ) Cite this article carbon steel rod

In this paper, the solid solution and precipitation behavior of inclusions on the surface and 1/2 thickness of the tested steel plate under the condition of welding heat input of 400 kJ/cm is investigated by using laser confocal experiments with hot-rolled state DH36 ship plate steel as the research object, and the mechanism of the effect of inclusions on the phase transformation of an acicular ferrite is revealed. The results show that the inclusions of the tested steel are mainly composed of Oxide-MnS, MnS, Oxide, TiN, Spinel, etc. The amount of inclusions on the surface of the tested steel plate is significantly higher than that at the 1/2 thickness position. During the heating stage, the small inclusions on the surface immediately disappeared, and the large inclusions gradually solidified in the matrix; atomic diffusion occurred at the bond between the inclusions and the matrix; while the small inclusions at the 1/2 thickness position gradually disappeared at the beginning of the heating stage, and the inclusions began to precipitate and grow when the temperature was increased to 990 °C. The acicular ferrite preferentially nucleates and grows near the boundary of the inclusions during the post-weld cooling stage, and its growth ends when two acicular ferrites cross.

Since high-heat input welding technology can greatly increase welding efficiency and shorten the construction time, it is increasingly used in offshore engineering1,2. The very short period of higher heat input to the heat-affected zone (HAZ) of the welded joint, especially the microstructure and properties of the coarse grain heat-affected zone (CGHAZ), will have a negative impact, ultimately affecting the low-temperature toughness of the tested steel welded joint and shortening the service life of the ship plate steel3,4.

By controlling the solid solution and precipitation behavior of non-metallic inclusions in the steel plate during the welding heat cycle, numerous researchers5,6,7, based on oxide metallurgy techniques, have generated diffusely distributed and fine (less than 3 μm in diameter) inclusions, suppressed the formation of CGHAZ tissue during high-heat input welding, and ultimately improved the mechanical properties of the welded joints of the tested steel. However, numerous types of inclusions, including oxides, sulfides, Mg-containing inclusions, TiN, and others, have different effects on the microstructure8,9. According to an argument made in10, oxide inclusions typically hurt the fatigue properties of a material, and low thermal expansion inclusions (such as Al2O3) produce residual stresses close to the yield strength of the matrix during the cooling phase; Shibata11 found that the addition of a moderate amount of Cr to Si-containing materials (˃1%) can induce the oxide inclusions MnO-SiO2 to MnO-Cr2O3, which is finer in size and has a strong pegging effect on grain boundary migration, thus inhibiting grain coarsening.

Sulfides (mostly MnS) are good nucleating agents for intracrystalline acicular ferrite (IAF), which can play a role in reducing interfacial energy and lattice mismatch12,13,14. It has been shown that MnS nucleates on MnO-SiO2, and then the Mn in the matrix diffuses to the MnS nucleus to make it grow; Kim15 discovered that MnS inclusion size and precipitation rate are both affected by the cooling rate, with faster cooling rate resulting in smaller MnS inclusion sizes and lower precipitation rates. Mg-containing inclusions (e.g., MgO-Al2O3-Ti2O3-MnS) can effectively promote IAF nucleation, and the addition of Mg elements can refine Ti-containing inclusions16; Xu17 found that lower Mg content (0.0027%) has little effect on the growth of austenite grains, and the large size of austenite grains is favorable for IAF nucleation; while a higher Mg content (0.0099%) inhibited the growth of austenite grains, which in turn inhibited the production of IAF. Wan18 found that fine TiN inclusions can nail austenite grain boundaries at high temperatures and limit grain growth, while the addition of Ti can reduce the interfacial energy between inclusions and matrix, promote MnS nucleation, form Mn-poor regions, and provide the driving force for IAF nucleation.

In previous research19, the authors discovered that the main control of inclusions in ship plate steel under high-heat input welding conditions can be divided into two stages: solid solution of inclusions during rapid temperature rise and precipitation of inclusions during post-weld cooling. The formation of fine-grained austenite is consistent with the phenomenon reported in the literature20,21,22, and may occur at the peak welding temperature when hard-to-solidify oxidized inclusions can peg the migration of grain boundaries of the austenite phase. The cooling process in the welding heat cycle can be used to improve the microstructure of the CGHAZ. This is achieved by introducing numerous fine and diffuse precipitates (micro- and nanoscale second phase) into the austenitic structure, which promotes IAF nucleation, growth and formation of the cross-interlocking structure of IAF. The increase in grain boundary area and the increase in grain boundary angle difference can effectively inhibit crack initiation and propagation, thus significantly improving the mechanical properties of the HAZ of welded joints23,24,25,26.

For the large size (thickness greater than 50 mm) of the ship plate steel, there are significant differences in the solid solution and precipitation behavior of inclusions at different locations in the thickness direction under high-heat input welding conditions. This area has been ignored in previous studies. Based on the calculation of the “2SEG Technical Note,” this paper demonstrates that the recommended heat input for high-heat input welding of a 60 mm thick steel plate is 400 kJ/cm. In this study, we conducted high-temperature laser confocal experiments to investigate the evolution of inclusions in the microstructure of steel under a welding heat input of 400 kJ/cm. The study focused on different locations in the steel thickness direction and aimed to reveal the solid-state phase change mechanism of inclusions on the microstructure.

The tested steel is hot-rolled DH36 grade ship plate steel, according to the DIN EN ISO 14284-2022 standard. The chemical composition of the tested steel was analyzed using titration techniques for accurate evaluation, and the results are shown in Table 1. The φ7 mm × 3 mm thermal simulation specimens were cut from the surface and 1/2 thickness of the tested steel plate using wire electrical discharge machining (EDM). The specimens were then ground, polished, and observed in situ using a high-temperature laser confocal experimental machine of the type VL2000DX Lasertec. The temperature interval for simulating the coarse crystal heat affected zone of high-heat input welded joints is usually set from 1350 to 500 ℃27,28. The process route is as follows: heating from 0 to 200 ℃ at a rate of 0.67 ℃/second to prevent temperature shocks and crucible rupture; then heating at the maximum rate of the equipment to a peak temperature of 1350 ℃, with a holding time of 1 s. Based on the t8/5 equation for welding heat input29, as shown in Eq. 1, the simulated welding heat input energy of 400 kJ/cm was calculated when the t8/5 was 176 s. Finally, in the stage from 500 °C to 0 ℃, the cooling rate of 120 ℃/min was used to simulate natural air cooling.

where E is the welding heat input energy, J; d is the plate thickness, mm; l is the thermal conductivity, W·m-1·K-1; ρ is the density, kg/m3; c is the specific heat capacity, J/kg·℃; t is the time, sec; T0 is the initial temperature, ℃. The experimental samples were ground and polished, etched with 4% nitric acid alcohol solution, and the microstructure was observed using an Axio Vert.A1 Zeiss microscope and ZEISS EVO Ma10 scanning electron microscope to observe the microstructure of CGHAZ in its coarse crystal heat-affected zone. The inclusions at different locations of the specimen steel plates before and after thermal simulation were also analyzed by OTS scanning, respectively.

The tested steel surface and 1/2 position before and after 400 kJ/cm heat input were analyzed by OTS for the inclusion size type and quantity, as shown in Fig. 1. Inclusions at the surface of the rolled tested steel were mainly Oxide-MnS and MnS with 25.2/mm2 and 18.2/mm2 respectively; after 400 kJ/cm heat input, the inclusions were mainly MnS and Oxide with 26.4/mm2 and 13.2/mm2 respectively; Overall, the number of units of Oxide-MnS, Sulfide, Spinel, TiN, Oxide-Sulfide and Silicate decreased with heat input in tested steel, as shown in Fig. 1a.

Type and quantity of inclusions (a)at the surface; (b)1/2 thickness.

The tested steel plate has inclusions mainly consisting of Oxide-MnS, MnS, Oxide, and TiN, with unit numbers of 5.8/mm2, 6.9/mm2, 5.9/mm2, and 7.1/mm2, respectively, at 1/2 thickness. After welding with a heat input of 400 kJ/cm, the main inclusions composition remains the same, with unit numbers of 5.8/mm2, 6.9/mm2, 5.9/mm2, and 7.1/mm2, respectively. However, the quantity per unit area of Oxide-MnS, MnS, Oxide, and TiN increases to 6.3/mm2, 18.0/mm2, 12.0/mm2, and 6.5/mm2, respectively, after the 400 kJ/cm welding heat input. The total number of inclusions per unit area also increases, as shown in Fig. 1b. Additionally, the number of inclusions at the 1/2 thickness position before and after heat input is lower than that at the surface position.

The inclusions accounted for more Oxide-MnS, Oxide, Spinel, and TiN. Based on the elemental distribution map and comparison with the results of previous studies17,18,30, the types of inclusions can be further confirmed, and the analysis results are shown in Fig. 2: (a) is Oxide-MnS composite inclusion; (b) is Oxide inclusion; (c) is Spinel inclusion; (d) is a typical rectangular TiN inclusion.

Energy spectrum and element distribution of inclusion.

Figure 3 counts the total number and size of inclusions at the surface and 1/2 of the tested steel before and after heat input. In Fig. 3a, the total number of inclusion particles at the surface of the original tested steel is 671 (74 /mm2); the size distribution of inclusions ranges from 0 μm to 7 μm, with the highest number of inclusions found in the size range of 1 μm to 2 μm, reaching 37/mm2. After treatment with 400 kJ/cm welding heat input, the total number of inclusion particles at the surface of the tested steel was 452 (55/mm2); the size of inclusions was mainly distributed between 0 μm and 6 μm. Figure 3b shows that there are 309 inclusions (34/mm2) on half the thickness of the original tested steel plate. These inclusions are identical to those at the surface of the original tested steel and are mostly in the 1 μm to 2 μm size range. The tested steel plate had a total of 308 inclusions at a thickness of 1/2 after the 400 kJ/cm welding heat input treatment (48/mm2), with most of them localized between 1 and 2 μm. This demonstrates that the number of inclusions at the surface position of the tested steel decreases significantly after welding heat input, while the number of inclusions at the 1/2 position increases.

Size and quantity of inclusions (a) at the surface; (b) 1/2 thickness.

Figure 4 shows the in-situ observation of inclusions on the surface of the tested steel with a welding heat input of 400 kJ/cm. As shown in Fig. 4, when the temperature reaches 875 ℃, the inclusions with a diameter of 4.2 µm undergo solid solution, and their progressive process becomes evident as the temperature rises. At a temperature of 1191 ℃, the inclusions are completely dissolved in the matrix. During the solid solution of the inclusions, the shadow range around them first increases and then decreases. Further observation reveals that the shadow becomes lighter the further away from the center of the inclusions, and darker the closer it is to the center. This is caused by the diffusion of atoms into the matrix.

In situ observation of inclusions at the surface of 400 kJ/cm.

Figure 5 presents the in-situ observation of the solid solution of the inclusions at the 1/2 position. The observations show that inclusions with dimensions of 3.8 μm and 4.5 μm dissolved as the temperature increased. The 3.8 μm inclusions began solid solution at 403.4 ℃ and completed the process at 463.5 ℃, while the 4.5 μm inclusions finished at 533.4 ℃. By the time the temperature reached 879.1 ℃, almost all inclusions within the field of view had undergone solid solution. Importantly, inclusions of similar size underwent solid solution without significant shadowing compared to the surface, indicating that they started and ended solid solution at lower temperatures than the surface.

In situ observation of inclusions at 1/2 thickness of 400 kJ/cm.

Figure 6 depicts inclusions that precipitated during a high-temperature process, as indicated by the black arrows and dashed rectangles. These inclusions start to form in the tested steel at 1254 ℃, with precipitation completing by 1327 ℃, defining this temperature range as the precipitation range for inclusions with high solid solubility. The figure also illustrates how the precipitated inclusions solidify back into the tested steel matrix. For example, the inclusions within dashed rectangle #1 in Fig. 6 grow at 1245 ℃, solidify again at 1311.8 ℃, and only a small residue remains at 1133.2 ℃. In dashed rectangle #2, there are three precipitated inclusions shown in Fig. 6, with the middle one starting to re-solidify at 1311.8 ℃, while the both sides inclusions persist at 1133.2 ℃, though the middle one vanishes. The figure also indicates that the inclusions grow at high temperatures, expanding from 4 μm at 993 ℃ to 7 μm at 1133 ℃, transitioning from irregularly round to uniformly round shapes. These inclusions are located at the austenite boundary and do not precipitate at the surface location under the same high-temperature conditions.

In situ observation of inclusion precipitation at 1/2 thickness of 400 kJ/cm.

The microstructure of the tested steel at the surface and half of its thickness is shown in Fig. 7. As depicted in the image, the microstructure mainly consists of polygonal ferrite, granular bainite, and a small amount of pearlite. The grain size of the polygonal ferrite was determined using the area method: the grain size at the surface was 8.4 μm and at 1/2 thickness it was 17.5 μm. This difference is due to the rolling process, where the deformation of larger core grains causes them to elongate, and the core of the thick plate has a slower cooling rate compared to the surface. This results in recrystallization growth during cooling. In general, the roll extrusion and flow have a significant impact on the hot rolling process, as shown in Fig. 8, indicating substantial deformation in the center of the steel plate. Consequently, the average grain size at the center of the steel is noticeably larger than at the surface, and the surface position is more prone to nucleate and precipitate inclusions.

Cross-sectional microstructure of DH36 steel (a) at the surface; (b) 1/2 thickness.

Schematic diagram of deformation of plate cross-section.

Additionally, Fig. 6 reveals that only half of the thickness shows precipitated inclusions under high-temperature conditions. This is a result of the high deformation in the core of the hot-rolled tested steel, storing more strain energy, and the high temperatures driving phase transformation, aiding inclusion precipitation. The surface exhibits finer grains without precipitated grain boundary positions, but at high temperatures, the austenite grain growth process occurs, with insoluble inclusions acting as grain boundary pinning agents. As a result of the synergistic effect of rolling and thermal simulation, larger surface inclusions are less susceptible to dissolution with increased heat input. Inclusions positioned at the half-thickness undergo a dissolution–precipitation transformation, leading to the formation of smaller inclusions within the grains. This microstructural evolution creates an environment conducive to the nucleation of IAF.

Figure 9 depicts the in situ observation of inclusions migrating along austenite grain boundaries during austenite growth at 400 kJ/cm heat input on the steel surface. The figure illustrates the process of austenite growth from 1335.9 °C to 1190.6 °C, showing the migration of the austenite grain boundary line L from left to right. At 1291.5 °C, inclusions are encountered, causing the grain boundary migration to halt. This results in the transformation of a large grain of austenite into two smaller grains on the left and right. The presence of fine inclusions aids in refining the grain boundaries of austenite and prevents the coarsening of grains in the heat-affected zone (HAZ).

Grain boundary migration of inclusion pinned austenite.

It is generally accepted in academia that the inclusions-induced acicular ferrite mechanism is related to the nature of the inclusions, mainly referring to the type of inclusions31,32,33. There are four main types of inclusions-induced acicular ferrite mechanisms, which are: the solute depletion in the vicinity of non-metallic inclusions, the low interfacial energy between inclusion and ferrite, the thermal strain energy due to the differential thermal contraction, the provision of an inert surface34.

Bamfitt35 proposed a two-dimensional model for calculating the degree of mismatch between inclusions and ferrite, and the formula is shown in Eq. 2.

where (hkl)s: refers to the ferrite low index crystal surface.

(hkl)n: refers to the inclusions low index crystal plane.

[uvw]s: refers to a low index crystal orientation on the (hkl)s crystal plane.

[uvw]n: refers to a low index crystal orientation on the (hkl)n crystal plane.

d[uvw]s: refers to the crystal plane spacing on the [uvw]s crystal plane.

d[uvw]n: refers to the crystal plane spacing from [uvw]n crystal up.

θ: refers to the angle between [uvw]s and [uvw]n.

The calculation results demonstrate that the inclusions promote IAF nucleation very effectively when the mismatch degree (δ) is 6%; at 6% < δ < 12%, they have some influence; and when δ > 12%, they have no effect. This type of inclusion readily triggers IAF nucleation because the misfit degree between MgAl2O4 and ferrite is 0.6%, which is much less than 6.0. MgAl2O4 forms a co-grid interface with ferrite to reduce the nucleation work of ferrite, and it is easy to induce IAF nucleation.

Figure 10 shows the element distribution of Spinel-type inclusions. As can be seen from the picture, the nucleation and growth of the IAF are induced by MgAl2O4 inclusion with a size of 5.8 μm. Barbaro’s research36 shows that when the inclusion size is greater than 0.4 μm, the IAF can be nucleated. According to the nucleation formula \({r}^{*}={\gamma }_{\alpha /\gamma }/\Delta {G}_{V}\) (where \({\gamma }_{\alpha /\gamma }\) is the energy of the ferrite/austenite phase boundary and \(\Delta {G}_{V}\) is the volume free energy of nucleus). The study12 shows that the key factor affecting nucleation is the volume free energy of critical nucleation of ferrite. However, some studies have demonstrated37,38 that different thermal shrinkage of shear stress generated between inclusions and matrix is also conducive to the growth of IAF.

Element distribution of MgAl2O4 inclusion.

Figure 11 illustrates the occurrence of inclusions-induced IAF at 400 kJ/cm heat input. As the temperature rises, the IAF expands to intersect with another IAF at approximately a 90° angle. This indicates that the generation of IAF can enhance grain refinement, promote interface growth, deter crack formation, and restrict crack propagation. The wide-angle grain boundary of IAF with a cross-interlocking structure can enhance the material's toughness.

Growth of IAF under 400 kJ/cm heat input.

The number of surface inclusions is significantly higher than those located at 1/2 the thickness of the rolled steel, primarily concentrated in the size range of 0–4 μm. After applying a welding heat input of 400 J/cm, many small-volume inclusions on the surface were dissolved into the matrix. In contrast, inclusions at 1/2 the thickness appeared to precipitate, with their quantity increasing after heat treatment. This precipitation is associated with deformation in the steel core during rolling and the storage of strain energy.

Because the tested steel surface is more sensitive to the influence of welding thermal cycle, it is susceptible to austenite grain growth under high-temperature conditions, which results in diminished welding properties. At this time, only the increased volume of insoluble inclusions acts as a barrier to restrict grain growth by pinning the grain boundaries. In addition, both before and after welding, the higher concentration of inclusions on the surface effectively promotes the nucleation of intergranular ferrite (IAF), resulting in a greater presence of IAF at the surface and maintaining a smaller austenite grain size.

All data included in this study are available upon request by contact with the corresponding author.

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This study was funded by the National Natural Science Foundation of China (No. 52004122). Liaoning Provincial University basic research funds project—Innovative talent cultivation special (Qihang Pang).

School of Materials and Metallurgy, University of Science and Technology Liaoning, Anshan, 114051, Liaoning, China

Jinming Shi, Qihang Pang, Weijuan Li, Zhaohua Xiang & Huan Qi

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S.J.M. and P.Q.H. wrote the main manuscript text. L.W.J. and Q.H. and X.S.H. prepared figures. All authors reviewed the manuscript.

The authors declare no competing interests.

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Shi, J., Pang, Q., Li, W. et al. Evolution of inclusions in DH36 grade ship plate steel during high heat input welding. Sci Rep 14, 18921 (2024). https://doi.org/10.1038/s41598-024-69907-1

DOI: https://doi.org/10.1038/s41598-024-69907-1

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Evolution of inclusions in DH36 grade ship plate steel during high heat input welding | Scientific Reports