JPS6239228B2 - - Google Patents

Description

[Detailed description of the invention]

The present invention relates to ultrafine-grained ferrite steel, particularly ultrafine-grained ferrite steel that is hot-rolled and is mainly composed of subeutectic steel that does not contain special alloying elements such as Nb, Ta, Mo, and W. It is. Various attempts have been made to refine the grains of ferritic steel materials. In other words, grain refinement in ferritic steel is the only way to simultaneously increase yield stress (also tensile strength) and improve toughness (fracture transition temperature). , a method of adding special alloying elements such as Nb and Mo, or a method of combining the two. However, a method of providing ultra-fine grained steel of 4 μm or less in an industrially viable manner with practical subeutectic steel has not yet been achieved. The present invention relates to such an epoch-making industrially obtained ultra-fine grain steel, which is characterized by weight percentages of C: 0.3% or less, Si: 1.5% or less, Mn: 2 % or less, N: 0.002 to 0.01%, balance Fe
It has high strength and ductility, and is characterized by containing 70% or more by volume of deformation-induced equiaxed ferrite crystal grains surrounded by large-angle grain boundaries with an average size of 4 μm or less in its metal structure. It is an excellent ultra-fine grained ferrite steel. This will be explained in more detail below. The steel of the present invention is an ultra-fine-grained ferrite steel as described above, but the structure called fine-grained ferrite in the present invention does not involve any significant elongation of the grain shape and is almost isotropic. refers to a structure consisting of crystal grains surrounded by so-called high-angle grain boundaries, and subgrain boundaries (low-angle grain boundaries) are not considered grain boundaries. The main reason for determining the composition range of the steel of the present invention is as follows. That is, although the carbon content was specified to be 0.3% or less, generally as the carbon content increases, the ferrite content inevitably decreases and the pearlite content increases. In the steel of the present invention, more ferrite is generated than expected from the normal phase diagram, but when carbon exceeds 0.3%, the amount of other structures such as pearlite increases, making it difficult to obtain a ferrite amount of 70% or more. Therefore, the above component range was set. Si is usually added to steel for the purpose of deoxidizing and is contained to some extent, and in the present invention, it is sometimes added intentionally because it has the effect of increasing the amount of ferrite. However, if it is added in excess of 1.5%, the ferrite crystal grains tend to become coarse, so 1.5%
The following was made. Furthermore, Mn adjusts the transformation point, making deformation-induced transformation more likely to occur, and contributes to grain refinement by preventing the rapid growth of deformation-induced ferrite, but if it is added in excess of 2%, the transformation temperature drops too much, resulting in ferrite. Since there are cases where the amount does not reach 70%, it is set to 2% or less. P and S are elements that are normally contained in steel to some extent, but if they are contained in large amounts, they impair the ductility of steel, but the amount contained in normal steel, P0.03%.
Hereinafter, the amount of S will not be particularly limited as it will not have a significant effect on the essence of the present invention if it is about 0.02% or less. Although some amount of N is contained in the steel as an impurity element, the amount is usually about 0.002 to 0.01%, and within this range it does not significantly affect the properties of the steel of the present invention. Note that when the amount of N is less than 0.002%, deformation-induced transformation occurs more easily than in the present invention, and when it exceeds 0.1%, it becomes less likely to occur, especially when elements such as Al and Ti are included. Al is usually contained in steel to some extent for deoxidation, but if it is normally contained at 0.1% or less, it generally does not have a large effect on the properties of the steel of the present invention. Further, Nb, Ta, Mo, W, etc. are all known as elements that delay the recrystallization and transformation of austenite. In the steel of the present invention, during hot working, the transformation of austenite into ferrite and/or the recrystallization of ferrite are promoted to make the grains finer, so Nb, Ta, Mo, W, etc. are elements that inhibit this. Therefore, it must not be included in the steel of the present invention. Steel with such a composition is manufactured by the following manufacturing method. The above steel may be continuously rolled at a total reduction rate of 50% or more in one or more passes in a temperature range that is substantially in the austenite region near the Ar ₃ transformation point of the steel, or in a short period of less than 1 second. For example, fine ferrite crystal grains are generated by rolling at a total reduction rate of 50% or more in one pass or two or more passes to cause transformation by rolling. To explain in more detail, the above steel has a normal Ar ₃ transformation point (the temperature at which the steel starts to undergo ferrite transformation during slow cooling from the temperature at which it is in the austenite region, hereinafter simply referred to as Ar ₃ ) and an Ar ₁ transformation point (similarly referred to as Ar 3). refers to the temperature at which pearlite transformation begins during slow cooling, and is simply referred to as
(Ar ₁ +50℃) ~ ( _{Ar 3}
+100°C) within 1 second with a total rolling reduction rate of 50% or more in 1 pass or 2 passes or more. The grain refining method according to the present invention is based on the principle that ferrite transformation is induced by strain applied to the austenite phase by rolling. Therefore, rolling, which is effective for grain refinement, is performed at a temperature below the equilibrium transformation point (Ae ₃ ), which is determined thermodynamically by the steel composition, and at which the normal austenite phase starts to transform during cooling (Ar ₃ ). From the above temperatures, the temperature will be around the Af ₃ point or not much lower than this. In normal transformation, the driving force for transformation to start and progress is supercooling, and in such transformation, the number of ferrite grains is mainly determined by the grain size of the austenite phase before transformation, and in particular, grain refinement measures are not taken. The ferrite grain size when not removed is usually about 8 to 10 microns or more. Processing strain also acts as a driving force for the initiation and progression of transformation, thus increasing the Ar ₃ point of austenite.
However, it has been known that ferrite transformation occurs immediately or after a very short period of time when austenite is subjected to a process such as rolling, and that the transformed ferrite grains are extremely fine. This is a new fact that has never been seen before, and the present invention is novel in that it thoroughly utilizes processing strain as a driving force for transformation. Fine ferrite grains precipitate at the grain boundaries of the austenite processed directly above the _three Ar points either immediately or after a short time after processing, and when further processing strain is applied, new ferrite grains precipitate at the interface between the ferrite grains and untransformed austenite. If this process is repeated and the processing strain is sufficiently large, it is thought that the entire surface will be covered with new ferrite grains at the end of rolling. At this time, the transformation nuclei generated during machining do not stop moving at the boundary of the transformation nuclei immediately after machining.
Also, since the processing strain (due to high density dislocations introduced by processing) does not recover immediately, the driving force against the movement of the interface is maintained to some extent, so fine-grained ferrite grows a little after processing, and its volume decreases. rate also increases. In such cases, it is difficult to experimentally strictly distinguish between the dynamically transformed ferrite part that has transformed during processing and the quasi-dynamic ferrite part that has grown following processing, and there is no practical need to do so. The state up to a time period of about 1 second or less in the problematic temperature range after processing is defined as processing-induced transformation. In addition, in the case of high-reduction machining even in one pass, and especially in the case of multi-pass machining of two or more passes, the ferrite generated in the first half of machining or before the final pass of the multi-pass continues to undergo machining, resulting in recovery of machining distortion. Therefore, recrystallization may occur. Actually 1
If the next process is added within a short time of about 1 second or less after completing a pass, the total reduction rate of these two passes will be reduced.
If it is larger than 50%, recrystallization nuclei will occur during machining, resulting in the condition in the first half of machining or after the previous pass machining.
In some cases, finer grained ferrite is produced compared to the state before the next pass processing. Even in such a case, it is difficult to distinguish recrystallization nuclei generated dynamically (during processing) immediately after processing from continuous growth (quasi-dynamic recrystallization) immediately after processing. Furthermore, the structure in which such dynamic and quasi-dynamic recrystallization phenomena are added is extremely similar to the above-mentioned structure in which only dynamic and quasi-dynamic transformation occurs, and it is difficult to clearly distinguish between the two. . Therefore, in the present invention, these are collectively referred to as deformation-induced ferrite. Such processing-induced ferrite is extremely fine and has the following characteristics. The shape is approximately equiaxed and no significant stretching is observed during processing. This will become clear from FIG. 1, which is a typical optical microscopic structure of the steel of the present invention (Test No. 4). This becomes clear when compared with Test No. 6 (Figure 2), which is a typical example of a so-called transformed material.
Test No. 6 shown in Figure 2 is a steel having the same composition as the invention steel (Test No. 4), but rolled at a lower reduction rate than the invention steel obtained at a slightly lower temperature range. This is an optical micrograph, and in this figure, the ferrite is stretched almost in the rolling direction (horizontal direction in the figure) and remains coarse as a whole. This type of structure is called a processed ferrite structure or a dynamic recovery structure, and when observed using a special corrosive solution, a structure called a sub-boundary is observed, which is different from the normal ferrite grain boundary. Unlike (high-angle grain boundaries), the difference in orientation between grain boundaries is small and the effect of grain boundaries on mechanical properties is completely different, so not only is there almost no improvement in properties due to fine grains, but also mechanical properties are improved by stretching. It exhibits significant anisotropy and is generally difficult to obtain the excellent properties of the steel of the present invention. Dislocation density is high on average and generally non-uniform. In the pro-eutectoid ferrite transformation that generally occurs during cooling, the transformation temperature is as high as 600 to 700°C, and the resulting ferrite is characterized by extremely few dislocations. On the other hand, as shown in FIG. 3, in the steel of the present invention (Test No. 4), there are areas where the dislocation density is locally high and areas where it is low, but overall it shows a considerably high dislocation density. This kind of dislocation density is observed in bainite or martensite, which has a low transformation temperature (about 500°C or less), and is extremely abnormal in the case shown in this figure, which is thought to have been formed at around 800°C. The reason for this is that dynamic transformation or dynamic recrystallization of ferrite that occurs during processing continues to be processed, so that the dislocation density increases as the processing progresses. For this reason, differences in dislocation density occur depending on the time of ferrite formation, and such non-uniformity takes on different aspects. Carbide is finely distributed almost uniformly. In the structure of the invention steel shown in Figure 1, the carbon content is 0.07%.
95% ferrite is produced, and no clear pearlite structure is observed. 0.12%C in Test 10 described below
98% ferrite has been obtained. The maximum amount of ferrite predicted from the equilibrium phase diagram is 91% and 85%, respectively, for these two steels, and the steel of the present invention produces far more ferrite than this. On the other hand, even if an optical microscopic structure shows only a very small amount of pearlite or bainite structure containing a large amount of carbide, or a martensitic structure containing a large amount of carbon as a solid solution, electron microscopic observation as shown in Figure 3 shows that it is fine. Fine precipitates (carbides) are observed at grain boundaries between grains or in areas with high dislocation density. In this way, carbides are precipitated finer and more uniformly than in ordinary ferrite pearlite steel. The reason for the formation of such an organization is as follows. As mentioned above, ultrafine-grained ferrite is produced near or above the normal transformation point, and in large amounts that would not be produced at that temperature. This is because they are supplied. On the other hand, as can be seen from the fact that a large amount of ferrite is produced within a short period of time, such as during rolling, carbon diffusion cannot keep up, unlike normal ferrite transformation. (In such a case, it is energetically disadvantageous compared to the case where transformation occurs due to carbon diffusion, but
This deficiency is also supplied through processing) Therefore, carbon atoms become a supersaturated solid solution in ferrite, and in the process of cooling to room temperature after processing, ferrite grain boundaries or high-density It is thought that this precipitates on the dislocations of The characteristics described above and in 2 are significantly different from ordinary pro-eutectoid ferrite. The fact that there is no diffusion of C is similar to the so-called Matsubu ferrite transformation observed in pure iron, etc., and the fact that the transformation progresses in an extremely short period of time during hot working is also consistent with this. The reason why the dislocation density is extremely high is that it undergoes processing at the same time as it is generated, and a steel with a new structure completely different from that of conventional steel, which can be called ``deformation-induced pine ferrite'', has been created. In addition, mechanical properties generally appear as the average of the properties of minute parts, so unless the ultra-fine grain structure occupies the majority, the properties will not be sufficiently exhibited. It is said that As mentioned above, one of the characteristics of the steel of the present invention is that depending on the manufacturing conditions, it is possible to generate a fine grain structure of nearly 100% regardless of the carbon content, and the higher this ratio, the more characteristic of the steel of the present invention. Needless to say, it appears strongly. However, as shown in Figure 4, steel with such a new structure exhibits properties that are particularly superior to those predicted by the conventional PETCH relational expression, but this is because the steel has a uniform structure as described above. This is thought to be due to the effect of ultra-fine graining. In addition to the characteristics shown in Figure 4, durability (fatigue strength),
It has been observed that it exhibits excellent properties in terms of formability (hole expandability, press formability), etc.
This all stems from the characteristics of the organization mentioned above. In addition, ultrafine grains of 4μ or more exhibit a superplastic effect in warm conditions. As mentioned above, the structure of the steel of the present invention is generated during or immediately after processing, so it is not necessarily necessary to increase the cooling rate after processing. Alternatively, in order to prevent grain growth during cooling that may occur in the case of low carbon, the cooling rate is generally 20℃/
It is desirable to cool the material within a range of sec or more to reach a temperature of 600°C or less. After that, various thermal histories can be taken depending on the required characteristics.

[Table] For example, to improve strength, it is sufficient to rapidly cool to around room temperature, and to improve workability,
After rapidly cooling to around 400°C, the solid solution carbon may be precipitated by slowly cooling from that temperature. Examples of the steel of the present invention will be described below. Thickness of converter melted steel containing the components shown in Table 1
Manufacture a 200mm CC slab and heat this slab to 1100℃
and hot rolled in a hot strip mill. Rough rolling was performed by rolling from 200 mm to 50 mm in 5 passes, and rolling was performed from 200 mm to 50 mm in 6 passes of finish rolling. A (the present invention) is an example in which a 58% reduction is applied within 1 second in the final two passes in finish rolling, and a 27% reduction is applied in the same final pass within 2 seconds.
The example in which the pressure was applied was designated as B (comparative example). The desired finishing temperature of the present invention is 680 to 870°C for A and B, based on the transformation points of the two types of steel.
The temperature is 660-890℃. Water cooling started about 1 second after the end of rolling. The finishing pass schedule for each
Shown in the table.

【table】

[Table] Table 3 shows the structure and mechanical properties of the above-mentioned steel types and steels of the present invention manufactured by the pass schedule.

[Table] In this table, test Nos. 1 to 4 and 8 to 10 are the steels of the present invention, and among them, test Nos. 1, 2, 4, 8, and 10 are the steels cooled to 600℃ or less after finish rolling. This is an example manufactured at a speed of 20°C/sec or more, and in comparative example Test No. 5, the rolling temperature was high (normal conditions), the amount of quenched structure of bainite and martensite was large, and the amount of ferrite was 40°C.
%, and has high strength but low ductility.
In test No. 6, the finish rolling temperature is low and the rolling is ferrite rolling, so the processed structure is a relatively coarse ferrite-pearlite stretched structure that contains sub-grains, but the average ferrite grain size is large, and therefore the ductility is somewhat insufficient. However, the strength is also decreasing. In addition, in Test Nos. 7 and 11, due to insufficient reduction, the ferrite grains were transformed during cooling, and the ferrite grains were not thin enough, and the second phase of pearlite and bainite accounted for about 40%, resulting in a slight increase in strength. Ductility is not sufficient. As described above, it is clear that all of the steels of the present invention have excellent properties in terms of strength-ductility balance, such as exhibiting far superior ductility at the same strength level compared to comparative steels with the same chemical composition. As already explained in Figures 1 to 3 above, which compare the microstructures of the inventive steel (Test No. 4) and the comparative steel (Test No. 6), the structural characteristics of the inventive steel are compared. It appears more clearly than steel. In the above explanation, the case where hot strips are manufactured as a product has been described, but it is clear that the steel material of the present invention can also be obtained by processes such as plate rolling, wire rod rolling, hot extrusion, etc. As detailed above, the steel of the present invention has at least 50%
It not only has a tensile strength of Kg/mm ² and a yield stress of 40 Kg/mm ² or more, but also has sufficient ductility and workability as a practical steel. , exhibits a superplastic phenomenon, and has significantly improved ductility and good friction bonding properties. Moreover, since a steel material having such characteristics can be obtained by hot rolling using a sub-euthogonal steel with a small amount of alloying elements, its industrial effects are enormous.

[Brief explanation of the drawing]

Figures 1 to 3 are micrographs of the metallographic structure, where Figure 1 is an optical microscope photograph of the steel of the present invention, Figure 2 is an optical microscope photograph of a comparative steel, and Figure 3 is an electron microscope photograph of the steel of the present invention. FIG. 4 is a diagram showing the relationship between grain size, yield stress, and toughness.

Claims (1)

[Claims] 1% by weight, C: 0.3% or less, Si: 1.5% or less,
Consisting of Mn: 2% or less, N: 0.002-0.01%, the remainder Fe and unavoidable impurities, and its metal structure contains deformation-induced equiaxed ferrite crystal grains surrounded by large-angle grain boundaries with an average size of 4 μm or less as a volume fraction. Ultra-fine grained ferrite steel with high strength and excellent ductility, characterized by a content of 70% or more.