JPH066748B2 - Manufacturing method for grain-oriented silicon steel sheet with excellent magnetic properties - Google Patents

Description

DETAILED DESCRIPTION OF THE INVENTION FIELD OF INDUSTRIAL APPLICATION The present invention relates to a method for producing grain-oriented silicon steel sheet having excellent magnetic properties.

As is well known in the prior art, grain-oriented silicon steel sheets are primarily used as iron cores for transformers and other electrical equipment, and are required to have excellent magnetic properties, such as magnetization characteristics and iron loss characteristics. Recently, with the advancement of silicon steel sheet manufacturing technology, the _B10 value (i.e., the magnetic field strength 10
It has become possible to obtain excellent products with a magnetic flux density exceeding 1.89 T (tesla), which is represented by the magnetic flux density in the rolling direction generated when the sheet thickness is 0.30.
mm unidirectional silicon _steel sheet _{, W 17/50 value} (i.e. magnetic flux density
Iron loss when magnetized at 1.7T and a frequency of 50Hz is 1.10W/K
It is now possible to obtain low iron loss coils of less than 1000 kJ.

As mentioned above, the basic requirement for obtaining silicon steel sheets with excellent magnetic properties is to obtain a (110)
It is necessary to fully develop the secondary recrystallized grains of [001] orientation. To achieve this, the (110)[001]
The presence of an inhibitor that strongly suppresses the growth of grains with undesirable crystal orientations other than the orientation, and the sharply aligned (1
10) It is known that the formation of a suitable primary recrystallization texture is necessary for the secondary recrystallization grains of the [001] orientation to fully develop.
Fine precipitates of Se, AlN, etc. are used, and grain boundary segregation elements such as Sb, As, Bi, and P are also used as needed.
It has also been practiced to use elements such as b and Sn in combination with the inhibitor to enhance the inhibitor's effect. On the other hand, to form an appropriate primary recrystallization texture, a method has been adopted in which the process conditions of hot rolling and cold rolling are appropriately combined, and for this purpose, a complicated process such as performing two cold rolling steps with an intermediate annealing in between has also been adopted.

Recently, there has been a trend toward converting the manufacturing method of silicon steel slabs, which are the raw material for manufacturing silicon steel sheets, from the conventional ingot-making and blooming method to the continuous casting method. However, when using such slabs manufactured by continuous casting, a new problem arises that did not occur in slabs manufactured by the conventional ingot-making and blooming method. Namely, the presence of MnS, MnS, which act effectively as inhibitors,
To obtain fine precipitates such as AlN, the slab must be heated to a high temperature of 1250°C or higher for a long time before hot rolling to fully dissociate and dissolve the inhibitor elements, and then the cooling process during hot rolling must be controlled to precipitate them to an appropriate fine size. However, in the case of continuously cast slabs, abnormal coarse growth of crystal grains is likely to occur while the slab is heated at high temperatures as described above,
These abnormally large grains can cause the formation of incompletely developed secondary recrystallized grains, known as band-like fine grain structures, in the silicon steel sheet, which can lead to deterioration of magnetic properties.

Several methods have already been proposed for preventing the formation of the band-like fine grain structure and improving the magnetic properties. For example, according to Japanese Patent Laid-Open Publication No. 55-119126, when a material slab is processed to a predetermined thickness by hot rolling, the structure immediately before recrystallization rolling is controlled to have 3% or more of γ phase precipitated in an α phase matrix, and this is then rolled at 1230 to 960°C.
The present inventors have also disclosed a method of performing high reduction rolling for recrystallization in a temperature range of 100°C to 30% or more per rolling pass.
The material slab is made to contain the required amount of C according to the amount of Si,
This paper discloses a method for effectively preventing the band-shaped fine grain structure from occurring in the finished product by generating a predetermined amount or more of the gamma phase in a specific temperature range during hot rolling, thereby splitting and destroying the crystal grains that have grown coarse during high-temperature heating of the raw material slab during the hot rolling process.

However, while the above-mentioned methods for producing a predetermined amount or more of the γ phase during hot rolling can prevent the band-like fine grain structure in the finished product, the desired magnetic properties may not always be sufficient. Moreover, the effect of preventing the band-like fine grain structure itself is very unstable, and in extreme cases, a fine grain structure may occur over the entire surface of the finished product, significantly deteriorating the magnetic properties. For example, there is a problem in that the stability most required for industrial production is lacking.

On the other hand, in recent years, methods have been developed to improve the primary recrystallization texture by effectively utilizing carbon or carbides contained in steel. For example, Japanese Patent Publication No. 14009/1963 discloses a method for improving the primary recrystallization texture by reducing the temperature of a hot rolled sheet before the first cold rolling from 790°C or higher to 500°C.
A method has been disclosed in which lens-shaped carbides of a size visible with an optical microscope (several μm) are precipitated within the crystal grains by rapidly quenching to a temperature of 40°C or less and then holding the material in the temperature range of 310 to 480°C. The relatively large carbides produced by this method are effective in dividing and refining the coarse hot-rolled elongated grains formed in the hot rolling process, and are not harmful to the development of secondary recrystallized grains.
It is believed that this method plays a role in eliminating the [011] oriented crystal grains at the desired stage of the cold rolling process. However, it has been difficult to sufficiently improve the magnetic properties using this method alone.

More recently, the solid solution C in the crystal grains has been
Alternatively, methods using fine carbides have been developed.
In Publication No. 182, AlN is used as an inhibitor,
The hot rolled sheet is annealed at high temperature and then rapidly cooled, and the final cold rolling reduction is
The method discloses that when a single intensive cold rolling pass of 80% or more is performed, at least one aging treatment is performed between cold rolling passes. In this case, the aging treatment requires holding at a temperature range of 50 to 350°C for 1 minute or more or holding at a temperature range of 300 to 600°C for 1 to 30 seconds, and it is said that carrying out the aging treatment multiple times is effective.

However, this method is uneconomical because it significantly reduces the cold rolling efficiency and increases the cost of heat treatment of the steel sheet. In addition, in Japanese Patent Publication No. 19377/1981, when AlN and Sb are added in combination as inhibitors, the temperature range for cooling after intermediate annealing is set to 700 to 900°C in order to fully exert the effect of this combination addition.
A method has been disclosed in which the steel sheet is slowly cooled in the range of 200 to 2000 seconds, and then immediately quenched to 200°C or less. However, if one were to try to achieve a process of slowly cooling the steel sheet between 700 and 900°C in 200 to 2000 seconds according to this method, it would be necessary to significantly modify the cooling zone of the continuous annealing furnace and provide a long slow cooling zone in which the steel sheet is heated and maintained at this temperature range, and it would also be necessary to operate continuously at an extremely low speed, which would result in a significant decrease in production efficiency and an increase in manufacturing costs, resulting in economical problems. Furthermore, it is difficult to use AlN or Al in combination with these methods.
Using a specific inhibitor called N-Sb,
The effect can only be achieved by combining a heavy cold rolling process of 100% or more.
The {111}〈11〉 orientation is significantly concentrated, and the (110)[0
01] orientation shows only weak accumulation as a secondary orientation, and (110)
This method is fundamentally different from the method of strongly concentrating the [001] orientation, and these methods could not be applied to the production of grain-oriented silicon steel sheets using MnS and MnSe, which have been commonly used as inhibitors in the past.

On the other hand, as one of the known methods for effectively utilizing carbon in steel to improve the texture within the range of the final cold rolling reduction suitable for S and Se as inhibitors, for example, Japanese Patent Publication No. 56-3892 discloses a method for reducing the temperature between 600 and 300°C to 150°C during cooling after intermediate annealing.
The method disclosed in this publication involves cooling the steel sheet at a cooling rate of 1/min or more and then aging it at the final cold rolling stage. In this case, the aging treatment is carried out at 100 to 400°C for 5 seconds to 30 minutes, and the aging treatment must be carried out at least once between cold rolling passes.
Therefore, this also leads to a decrease in cold rolling efficiency and an increase in heat treatment costs, which is economically disadvantageous, and therefore there has been a strong demand for the development of a more efficient method.

Problems to be Solved by the Invention As mentioned above, the conventional methods for effectively utilizing the carbon in steel have had problems such as not being able to obtain sufficient magnetic properties or requiring special processes such as slow cooling at high temperatures or long-term aging treatment, which are economically disadvantageous.

This invention has been made in view of the above circumstances, and
By eliminating and improving the drawbacks of conventional methods for making effective use of
The object of the present invention is to provide a method for efficiently and economically producing grain-oriented silicon steel sheets having excellent magnetic properties on an industrial scale.

Means for Solving the Problems The present inventors have conducted extensive experiments and studies in order to achieve the above-mentioned object, and as a result have found that a grain-oriented silicon steel sheet having excellent magnetic properties can be produced efficiently and economically by combining the following three requirements: first, adjusting the C content in accordance with the Si content so as to control the amount of γ phase formed during hot rolling within an appropriate range; second, decarburizing a predetermined amount of C during the period from the end of the hot rolling process until after intermediate annealing before final cold rolling; and third, performing a process after intermediate annealing before final cold rolling to control and sufficiently disperse carbides within the crystal grains of the steel sheet within a specific, extremely small range. This discovery led to the present invention.

Specifically, the method for producing grain-oriented silicon steel sheets of the present invention is to use a steel sheet containing 0.015 to 0.10% of C, 2.8 to 4.0% of Si, and 0.02 to 0.15% of Mn.
and one or both of S and Se in a total amount of 0.008 to 0.080%, with the balance being substantially Fe, is hot-rolled, and the obtained hot-rolled steel sheet is cold-rolled two or more times with intermediate annealing in between, with a final cold reduction of 40 to 100%.
In a series of manufacturing methods for grain oriented silicon steel sheet, the C content of the silicon steel material is set within the range expressed by the following formula according to the Si content, and C is decarburized by 0.006 to 0.020% between the end of the hot rolling and before the end of the final cold rolling, and in the cooling process after intermediate annealing before the final cold rolling,
Rapid cooling within 30 seconds from 770 to 100°C, and immediately reduce to 150°C.
Aging treatment is performed at a temperature of 250°C for 2 to 60 seconds, or
Alternatively, during the cooling process after the intermediate annealing,
Rapid cooling in the temperature range of 300°C within 20 seconds, followed by 300 to 15
This method is characterized by controlling the cooling time required between 0°C and 100°C within a range of 8 to 30 seconds, thereby controlling the carbides within the grains of the steel sheet to a fine and well-dispersed precipitate state of 100 to 500 Å in size, and then performing final cold rolling.

0.37[Si%]+0.27≦log([C%]×10 ³ )≦0.37
[Si%] + 0.57 where [Si%] and [C%] are the S contained in the steel.
i represents the weight percent of C.

First, the findings that led to the development of this invention will be explained. The present inventors have investigated the effects of the γ phase formed during hot rolling and have found the following facts.
That is, the present inventors have newly discovered the following facts: the gamma phase formed during the hot rolling of a material slab is effective in splitting and destroying the crystal grains that have grown coarsely when the material slab is heated to high temperatures, as described above; however, it also has a detrimental effect on fine precipitates such as MnS and MnSe, which act as inhibitors; and in particular, excessive gamma phase formation may significantly reduce the inhibitor effect and hinder the sufficient development of secondary recrystallized grains. Therefore, the amount of gamma phase formed must be kept within an appropriate range; and even if the amount of gamma phase formed is within an appropriate range, after the gamma phase has played a role in subdividing the coarsely grown grains during hot rolling, it becomes detrimental to the formation of an appropriate crystal structure and texture in the cold rolling process. Therefore, the inventors have conducted various studies into measures to eliminate the harmful effects of the γ phase while utilizing its beneficial effects. As a result, they have found that by adjusting the amount of C in the material according to the amount of Si so that the amount of γ phase formed during hot rolling falls within an appropriate range, and by carrying out an appropriate amount of decarburization between the end of hot rolling and before the end of the final cold rolling process to reduce the amount of excess γ phase formed, and by controlling and sufficiently precipitating and dispersing the intragranular carbides in the steel sheet after intermediate annealing and before final cold rolling to within a specific, extremely small range that has not been noticed before and that is so small that it cannot be seen with an optical microscope, the texture of the steel sheet after final cold rolling and decarburization annealing but before final annealing can be improved to a state in which the concentration of the (110)[001] orientation is strong, and as a result, a highly uniform (1
10) We have newly discovered that grain-oriented silicon steel sheets with excellent magnetic properties can be obtained by sufficiently growing secondary recrystallized grains in the [001] orientation, which has led to the completion of this invention.

Based on the experimental results of the inventors who have completed the present invention as described above, each of the requirements of the present invention will now be described in more detail.

Figure 1 shows the results of the experiment using 0.015-0.035% Se as an inhibitor and 0.035% M
The content of Si is 2.8 to 3.1%, and the content of
The test pieces were divided into three groups: 1.0% to 3.5%, 3.6% to 3.8%, and the C content was varied within the range of 0.01% to 0.10%, with the balance being essentially Fe. A large number of silicon steel continuous cast slab test pieces, each with a composition consisting essentially of Fe, were heat treated at 1400°C for one hour and then hot rolled to a thickness of 2.
The sheet is then hot-rolled to a thickness of 5 mm, and then cold-rolled twice with intermediate annealing by a known method to a final thickness of 0.30 mm.
Furthermore, the iron loss _{W17 / 50} of each grain-oriented silicon steel sheet product obtained by decarburization annealing and final annealing was investigated, and the relationship between the iron loss value and the Si content and C content of each continuously cast slab test material is shown. In this test, the intermediate annealing atmosphere was changed from decarburizing to non-decarburizing, and the final cold rolling reduction was set in the range of 50 to 70%.

The symbols ◎, ○, ●, and × in Figure 1 indicate the iron loss W ₁₇ of the product.
The magnitude of the _{/ 50} was judged according to the Si content of each test material stage, as shown in Table 1 below.

The dashed lines A, B, C, D, and E shown in Figure 1 are estimated values for the amount of gamma phase formed at 1,150°C during hot rolling, and represent cases where the amount of gamma phase formed is 40%, 30%, 20%, 10%, and 0%, respectively. The amount of gamma phase formed substantially varies depending on the Si content, the C content, and the temperature, and the dashed lines A, B, C, D, and E were calculated from the following equation (1), which was derived from the correlation between the Si content and the C content in the steel and the actual measured value of the amount of gamma phase formed in an equilibrium state at 1,150°C, which was determined by experiments using silicon steel specimens with various Si and C contents.

γ%=67×log([C%]×10 ³ )−25[Si%]−8
...(1) As is clear from Fig. 1 and Table 1, the absolute iron loss level judged as good varies depending on the Si content. However, the appropriate range of the C content for excellent iron loss S _{17 / 50} according to each Si content is between the broken lines D and B, that is, when the amount of γ phase generation is 10
It was found that the effect was limited to within the range of ∼30%.
However, the γ phase generated during the hot rolling process is metastable, unlike the equilibrium state.
It is difficult to accurately grasp the amount of γ
Since it is not practical to limit the amount of phase produced,
It is considered appropriate to limit the range of C content according to the Si content in the material so that the estimated γ phase formation amount given by the above formula (1) is within the range of 10 to 30%. Based on this idea, in this invention, the range of C content according to the Si content in the material so that γ% is 10 to 30% is derived from the above formula (1), and this is set as the appropriate range of C content to obtain an excellent iron loss level. In other words, this appropriate range of C content is given by the following (2)
It is expressed by the formula:

0.37[Si%]+0.27≦log([C%]×10 ³ )≦0.37
[Si %]+0.57 (2) This is the first characteristic requirement of the present invention.

When the C content is less than the lower limit of the range of the appropriate C content according to the Si content shown in the above formula (2), that is, when the composition corresponds to the amount of γ phase generated during hot rolling being less than 10%, the crystal structure of the product shows a clear band-like fine grain structure, and deterioration of the magnetic properties is observed.
It was found that products with a composition of 10% or more, as shown by the line, hardly have any band-like fine grains, and are mostly composed of normally developed secondary recrystallized grains. Therefore, in order to split and destroy the coarse crystal grains that grow abnormally during high-temperature heating of the slab during the hot rolling process and prevent the generation of band-like fine grains in the product, it is necessary to generate a certain amount of gamma phase or more. The required amount of gamma phase is 10% if equilibrium is reached during hot rolling depending on the amount of Si contained.
It was found that this could be achieved by adding an amount of C that would produce the above-mentioned gamma phase. On the other hand, when the C content was significantly excessive, i.e., when the composition corresponded to one in which the amount of gamma phase produced during hot rolling exceeded 30%, the crystalline structure of the product became a fine-grained structure with incomplete secondary recrystallization, and the product exhibited extremely poor magnetic properties.

As described above, only when the C content is such that a gamma phase of 10 to 30% is formed in equilibrium during hot rolling, depending on the Si content, can the occurrence of band-like fine grain structures in the product or the formation of a fine grain structure over the entire surface with incomplete development of secondary recrystallized grains be prevented. Therefore, it has been found that limiting the C content in accordance with the Si content using the above formula (2) is extremely effective in improving magnetic properties.

However, even within the range of 10-30% of the gamma phase formation amount shown in Figure 1, there are still some steels with insufficient iron loss characteristics, and from the viewpoint of industrial production where stable magnetic properties must be ensured, it cannot be said that the restriction of the C and Si contents according to the above formula (2) alone is satisfactory. Therefore, the present inventors have conducted further research to improve this, and as a result have found that decarburizing the C by 0.006-0.020% during the process from the end of the hot rolling process of the material slab until after the intermediate annealing before the final cold rolling process is effective for consistently obtaining excellent magnetic properties, and this is the second characteristic requirement of the present invention.

This requirement was clarified by the inventors' experiments as follows. Specifically, among the test materials used in the experiment shown in Figure 1, those with two Si contents, 2.8-3.1% and 3.3-3.5%, and whose C contents corresponding to these Si contents were within the range corresponding to 10-30% of the γ phase formation at 1150°C during hot rolling, were used. The relationship between the magnetic properties of the products and the difference in C content immediately after the hot rolling process and after intermediate annealing before final cold rolling, i.e., the amount of decarburization ΔC between them, was investigated in detail, and the results shown in Figures 2(A) and (B) were obtained. In Figure 2, the white circles indicate the Si contents
2(A) and (B), it was found that excellent magnetic properties were stably obtained when the amount of decarburization ΔC was 0.006% or more and 0.020% or less, but when C was less than 0.006% or more than 0.020%, the magnetic flux density was insufficient and the iron loss was large, making it impossible to obtain sufficient magnetic properties.

In the ordinary production of silicon steel sheet, the amount of decarburization between hot rolling and final cold rolling is about 0.005% or less, and therefore the decarburization amount of 0.006-0.020% in the method of the present invention is larger than the amount of decarburization in the ordinary method. Therefore, in carrying out the method of the present invention, it is usually necessary to carry out an active decarburization treatment, such as by making the intermediate annealing atmosphere decarburizing. By carrying out an appropriate amount of strong decarburization between the end of hot rolling and before final cold rolling, it is possible to compensate for the shortcomings of the first requirement explained above and to stably obtain excellent magnetic properties.

As mentioned above, the fact that an appropriate amount of decarburization is effective in improving and stabilizing magnetic properties is also clear from the following observation results of the crystal structure and texture. That is, when the amount of decarburization is appropriate,
The grain size before final cold rolling is uniform and appropriate, and the primary recrystallization texture is improved to a favorable state showing a strong concentration of (110)[001] orientation. As a result, the crystalline structure of the product has well-developed normal secondary recrystallization grains.
On the other hand, when the amount of decarburization is insufficient, the primary recrystallization structure has irregular grains and residual clumps of carbides, and the primary recrystallization texture is inappropriate, with weak concentration of the (110)[001] orientation and dispersed {111}<11> orientation, resulting in poor secondary recrystallization development with a mixture of fine grains. Furthermore, when decarburization is excessive, the grain size before final cold rolling is inappropriate, with dispersed coarse grains. The primary recrystallization texture also has a reduced (110)[001] orientation, resulting in significantly coarse grains after secondary recrystallization, many of which are slightly deviated from the (110)[001] orientation, resulting in insufficient magnetic properties.

As mentioned above, the inventors have found that an appropriate amount of decarburization is effective in improving and stabilizing magnetic properties. Furthermore, the inventors have found that a higher magnetic flux density and iron loss of _1.00 W _{/ 50} are obtained.
As a result of efforts to develop grain-oriented silicon steel sheet with extremely excellent properties of less than 1000000000 kg, it was discovered that by combining the above two requirements with a process for controlling intragranular carbides in the steel sheet to a very small, specific range that is not visible with an optical microscope and precipitating a sufficiently large amount after intermediate annealing before final cold rolling, it is possible to improve the texture before final annealing to a state in which the concentration of (110)[001] orientation is even stronger, and as a result, highly uniform secondary recrystallized grains with (110)[001] orientation are formed in the secondary recrystallization process during final annealing, and excellent magnetic properties are obtained. This process for controlling intragranular carbides is defined as the third characteristic requirement of this invention.

The effect of the third requirement will be explained below based on the results of experiments conducted by the inventors. The material used in the experiments was C0.045%, Si
The steel sheet has a composition containing 3.20%, 0.06% Mn, 0.030% Se, and the balance being substantially Fe. The steel sheet is manufactured through ordinary steelmaking, continuous casting and hot rolling into a plate having a thickness of 3.0 mm.
This hot-rolled sheet is annealed at 950°C for 2 minutes, pickled and subjected to the first cold rolling, resulting in an intermediate sheet thickness of 0.
After cutting to 75 mm, intermediate annealing was performed at 900°C for 3 minutes, and then the reduction rate was 60
The steel was then decarburized in a wet hydrogen atmosphere at 800°C, coated with MgO, and subjected to final annealing at 1200°C for 10 hours to obtain a grain-oriented silicon steel sheet product.

In the above experiment, the amount of decarburization ΔC during intermediate annealing between cold rolling processes
was changed to three levels: 0.002%, which is the conventional normal level, 0.012%, which is within the limited range of this invention, and 0.025%, which is excessive decarburization.
The following cooling is performed by oil quenching (cooling time of approximately 770 to 100°C).
The steel sheets were then subjected to a rapid cooling (equivalent to a quenching time of 10 seconds) and immediately aged at 200°C for periods varying from 2 to 200 seconds. Figure 3 shows the relationship between the size of intragranular carbide precipitates and magnetic properties of the aged steel sheets, i.e., steel sheets after intermediate annealing and before final cold rolling, as well as the relationship between carbide precipitate size and aging time at 200°C. The magnetic property plot in Figure 3 indicates a decarburization amount ΔC of 0.002%, a ●, and a ◎, respectively. The comparative material in Figure 3 is a sample that was forced air-cooled at a cooling rate equivalent to the 98-second cooling time between 770°C and 100°C, which is commonly used in industrial continuous annealing.

As is clear from FIG. 3, when the amount of decarburization is an appropriate amount (marked with a black circle) within the range of the second requirement and the aging treatment time at 200°C is about 10 to 20 seconds, the magnetic flux density _B10 value is
The specimen exhibited extremely excellent magnetic properties, with a magnetic field strength of 1.94 T or more and an iron loss _{W17 / 50} of 1.00 W/Kg or less, and it was clear that the size of the carbide precipitates in this case was in the range of 100 to 500 Å. An electron microscope photograph (10,000 magnifications) of the carbide precipitates in this case is shown in Figure 4(A). However, this electron microscope photograph was taken of a specimen that had been subjected to intermediate annealing before final cold rolling, quenched between 770 and 100°C for 22 seconds, and then immediately aged at 200°C for 10 seconds. The average carbide grain size was 200 Å.
It is clear that the carbides are dispersed uniformly and in large quantities.

On the other hand, in the case of the specimens oil-quenched after intermediate annealing (without aging) and aged at 200°C for 2 seconds, the magnetic properties were clearly insufficient for both decarburization amounts. In these cases, intragranular carbides were not observed or only small amounts were locally precipitated. Furthermore, in the case of specimens aged at 200°C for 30 seconds or longer, the magnetic properties were clearly insufficient for both decarburization amounts. In this case, the intragranular carbide precipitate size exceeded 500 Å. For reference, Figure 4(B) shows an electron microscope photograph (10,000x magnification) of the carbide precipitate state before final cold rolling for a comparative specimen that underwent industrial standard cooling (approximately 98 seconds between 770°C and 100°C) after intermediate annealing. In this case, the average diameter of the intragranular carbide precipitates was approximately 700 Å, and the magnetic properties were significantly improved by quenching at 200°C after intermediate annealing.
The results were as poor as when the aging treatment was carried out at 200°C for 30 seconds or more.

Furthermore, from Figure 3, it is clear that when the amount of decarburization ΔC is at the conventional normal level (marked with ○) or when the amount of decarburization is excessive (marked with ◎), even when the steel is quenched after intermediate annealing and immediately subjected to aging treatment at 200°C for about 10 to 20 seconds, some improvement in magnetic properties is observed, but it is not significant.

From the above experimental results, it was found that the magnetic properties can be significantly improved by performing a process to set the intragranular carbide size within the range of 100 to 500 Å after intermediate annealing and before final cold rolling, particularly for materials with an appropriate amount of decarburization.

Furthermore, the inventors have found that (A) decarburization is not actively performed in the intermediate annealing process, and that standard cooling (a required cooling time of about 90 minutes between 770 and 100°C) is performed without rapid cooling in the cooling process after the intermediate annealing before final cold rolling.
(B) When 0.006 to 0.020% decarburization is performed in the intermediate annealing process and standard cooling is performed without rapid cooling in the cooling process after intermediate annealing before final cold rolling. (C) When decarburization is not actively performed in the intermediate annealing process and standard cooling is performed without rapid cooling in the cooling process after intermediate annealing before final cold rolling.
The temperature range is quenched within 30 seconds, and then immediately cooled to 200°C for 10-20 seconds.
When the cold-rolled sheets obtained by the above four types of treatments (A) to (D) were subjected to aging treatment of about 1.5 seconds, (D) decarburization of 0.006 to 0.020% was performed in the intermediate annealing process, and quenching and aging treatment similar to (C) was performed in the cooling process after the intermediate annealing before final cold rolling, the Goss orientation strength of the surface layer of the decarburized annealed sheet before final annealing was examined, and the results shown in Figure 5 were obtained. From Figure 5, it can be seen that when neither decarburization nor quenching temporary aging treatment was performed (A), the Goss orientation strength of the surface layer of the decarburized annealed sheet (B) and the quenching temporary aging treatment (C) were about 1.5 seconds.
The Goss orientation strength is doubled, and when both decarburization and quenching temporary aging treatment are performed according to the method of this invention,
It was confirmed that (D) exhibited approximately 1.5 times the Goss orientation strength compared to (A). The reason why the Goss orientation strength is increased by the method of this invention is thought to be as follows: An appropriate amount of decarburization lowers the recrystallization start temperature during intermediate annealing before final cold rolling, which is advantageous for the growth of Goss grains, which are said to recrystallize at lower temperatures. Furthermore, the reduction in the amount of α-γ transformation during soaking after recrystallization prevents the randomization of the texture, resulting in an improved texture with a strong concentration of Goss orientation. Furthermore, the uniform precipitation and dispersion of ultrafine carbides before final cold rolling serves to amplify the difference in the amount of internal strain accumulated during final cold rolling, which depends on the initial crystal orientation. It is presumed that, when recrystallization occurs during the subsequent temperature-raising process of decarburization annealing, crystal grains with the (110)[001] orientation and its neighboring crystal orientations, which have accumulated a greater amount of strain inside the crystal after cold rolling, will preferentially initiate recrystallization at an earlier stage, forming a primary recrystallized structure with a stronger Goss orientation. Therefore, the method of the present invention achieves an improved texture with a stronger concentration of Goss orientation through the synergistic effect of the above two actions.

On the other hand, if the amount of decarburization before the final cold rolling is insufficient, the primary recrystallization structure before the final cold rolling has an uneven grain size, with fine grains distributed in a blocky shape, and the primary recrystallization texture is (110)[001]
The orientation is weakly concentrated, and the relatively strong (111) <11> orientation is dispersed, resulting in an inappropriate structure. Even if rapid cooling is performed before final cold rolling to uniformly precipitate and disperse fine carbides of 100 to 500 Å, this has little effect, and as a result, the product's crystal structure is in a state of poor secondary recrystallization, with fine grains mixed in.

In addition, if the decarburization is excessive, the grain size before the final cold rolling becomes inappropriate with uneven and coarse grains dispersed, and the primary recrystallization texture also has a reduced (110)[001] orientation. In addition, excessive decarburization results in an insufficient amount of carbide precipitation during cooling in the intermediate annealing before the final cold rolling, and the amount of the desired fine carbide cannot be secured by rapid cooling. Therefore, the crystal structure of the product obtained in this state is dominated by extremely coarse secondary recrystallization grains, and these coarse grains are in the (110)
Many of the orientations are slightly deviated from the [001] orientation, which results in insufficient magnetic properties and a tendency for iron loss values to increase.

As described above in detail, only when an appropriate amount of decarburization before final cold rolling is combined with the desired intragranular carbide size can an extremely low iron loss value and a sufficiently high magnetic flux density be obtained. Even if the amount of decarburization is within an appropriate range, if intragranular carbides do not precipitate or grow to exceed 500 Å, or conversely, even if the intragranular carbide precipitate size is within the range of 100 to 500 Å, if the amount of decarburization before final cold rolling is too much or too little, the desired magnetic properties cannot be obtained.

Next, a specific method for sufficiently precipitating ultrafine carbides in the range of 100 to 500 Å within the crystal grains before the final cold rolling as described above will be explained.

Figure 6 shows the cooling time required for 22 minutes between 770 and 100°C after intermediate annealing.
The figure shows the relationship between the aging temperature and treatment time and the size of intragranular carbide precipitates when the specimen is quenched in 20 seconds and then immediately aged in the temperature range of 100 to 300°C. From Figure 6, in order to precipitate ultrafine carbides in the range of 100 to 500 Å by aging treatment after quenching, the specimen is required to be aged in the temperature range of 150 to 250°C for 2 to 60 seconds.
However, it was found that the lower the temperature, the more appropriate it is to select a longer holding time. Here, during cooling after intermediate annealing before final cold rolling, the amount of solid solution of C reaches a maximum at 770°C, so if the cooling rate in the region below 770°C is slow, coarse carbides will precipitate at grain boundaries and the like before the start of precipitation of fine carbides, and the desired amount of precipitation dispersion of fine carbides will not be obtained, making it impossible to improve the texture. Therefore, the cooling before aging treatment was decided to be rapid cooling between 770 and 100°C for 30 seconds or less.

Furthermore, the present inventors have attempted to develop a method for eliminating the need for aging treatment after cooling by strictly controlling the cooling process after intermediate annealing, particularly the cooling process below 300°C, which is a temperature range that has been overlooked in the past. That is, as can be seen from Figure 6, focusing on the fact that ultrafine carbides precipitate within grains in the temperature range of 300°C or less and 150°C or more, the inventors have attempted to develop a method for eliminating the need for aging treatment after cooling by strictly controlling the cooling process below 300°C, which is a temperature range that has been overlooked in the past.
The steel was rapidly cooled in the temperature range of 300 to 150°C in the same manner as above, and then cooled at various cooling rates in the temperature range of 300 to 150°C, and attempts were made to precipitate intragranular ultrafine carbides during the cooling between 300 and 150°C. Specifically, during the cooling after intermediate annealing before final cold rolling, the steel was cooled at various cooling rates between 770 and 300°C.
After quenching in the 150°C range with mist jet cooling for a required cooling time of 15 seconds, the steel was subsequently cooled in the temperature range below 300°C at various cooling rates, from water cooling to natural cooling, and the relationship between the cooling time between 300 and 150°C, the size of intragranular carbide precipitates, and the magnetic properties of the product was investigated, with the results shown in Figure 7. However, the amount of decarburization in the intermediate annealing before final cold rolling was 0.012%, which is within the range of this invention.

From FIG. 7, it is clear that in order to obtain intragranular carbide precipitation sizes of 100 to 500 Å, the cooling time between 300 and 150°C should be selected within the range of 8 to 30 seconds, and it is also clear that in this case, a significantly low iron loss value and a sufficiently high magnetic flux density can be obtained.

As described above, the grain size of the steel sheet before final cold rolling is 100 to 500.
As an industrial method for dispersing and precipitating ultrafine carbides of Å size, the steel is rapidly cooled between 770 and 100°C within 30 seconds during the cooling process of intermediate annealing before the final annealing.
It has been found that the following methods are suitable: aging at 150-250°C for 2-60 seconds; or rapid cooling between 770-300°C in 20 seconds or less, followed by controlling the cooling time between 300-150°C within a range of 8-30 seconds. Both of these methods can be easily implemented industrially, but the latter method has the advantage of enabling efficient continuous furnace operation by shortening the cooling time.

Next, the reasons for limiting the chemical composition of the silicon steel material used in the method of the present invention will be explained.

Silicon is an element effective in increasing resistivity and reducing iron loss, and if it is less than 2.8%, it is not possible to achieve a sufficiently low iron loss value, while if it exceeds 4.0%, it becomes extremely brittle and the cold rolling workability deteriorates, making ordinary industrial cold rolling difficult, so the range of silicon is limited to 2.8 to 4.0%.
Generally, the higher the content within the range of S, the lower the iron loss of the product can be obtained.
Increasing the amount of silicon not only increases the cost of silicon raw materials but also leads to increased costs due to a decrease in cold rolling yield.
The i content must be appropriately selected depending on the desired iron loss level to be obtained.

As mentioned above, the C content should be adjusted to fall within the range of the formula (2) in accordance with the Si content. That is, the C content must be in the range corresponding to the amount of γ phase formation of approximately 10 to 30% at 1150°C during hot rolling as shown in Figure 1. Specific numerical values according to the formula (2) are shown in the following Table 2.

However, if the C content is less than 0.015%, the required amount of γ phase cannot be secured when the Si content is in the range of 2.8 to 4.0%.
If the C content exceeds this range, the decarburization process will take a long time, which will be economically disadvantageous. Therefore, it is necessary to satisfy the above formula (2) within the range of 0.015 to 0.10%.

Mn, S, and Se are all added as inhibitors to suppress the growth of primary recrystallized grains during final annealing.
10) It is an element necessary for the sharp development of secondary recrystallized grains of the [001] orientation. However, Mn 0.02 to 0.15%, S, S
One or two of the above e in total amount of 0.008 to 0.080%
If the content is outside the above range, either too much or too little, the secondary recrystallization becomes unstable and the desired excellent magnetic properties cannot be obtained, so the content is limited to the above range.

The silicon steel material to which the method of the present invention is applied consists essentially of Fe and unavoidable impurities in addition to the above-mentioned components.

Next, the entire manufacturing process of the grain-oriented silicon steel sheet according to the method of the present invention will be explained in order of steps.

The silicon steel slabs used in this invention may be those obtained by the conventional ingot-making and blooming process or by continuous casting. However, the method of this invention is particularly effective in stabilizing and improving magnetic properties when using continuously cast slabs. In this method, the silicon steel slab is heated to about 1250°C or higher, hot rolled by a known method, and finished into a hot-rolled sheet with a thickness of 1.2 to 5.0 mm. This is then subjected to normalizing annealing at 750 to 1100°C as needed, and then cold-rolled two or more times with intermediate annealing at 750 to 1100°C between them to obtain a final cold-rolled sheet with a final thickness of 0.15 to 0.50 mm. During this process, between the steps after hot rolling and before final cold rolling, i.e., during the self-annealing after hot rolling and coiling or at least one of the normalizing annealing and intermediate annealing steps, the atmosphere is adjusted to a decarburizing level, thereby achieving a total decarburization of 0.006 to 0.020%. The strength of the decarburization in the decarburization annealing atmosphere should be adjusted appropriately depending on the material composition, plate thickness, annealing time, etc. Furthermore, when using the self-annealing time after coiling of the hot rolled coil, it is necessary to add Fe ₂ O between the coil layers.
It is also possible to perform decarburization annealing of the hot-rolled sheet by applying an oxide such as _Cr-3 .

In the cooling process of intermediate annealing before final cold rolling in the cold rolling process, the above-mentioned cooling methods are used to sufficiently precipitate ultrafine carbides of 100 to 500 Å in size within the crystal grains of the steel sheet before final cold rolling, and then the final cold rolling reduction ratio is 40
In this invention, the crystalline structure is homogenized by performing appropriate decarburization and fine precipitation of carbides before final cold rolling, and a strong concentration of the (110)[001] orientation in the texture is promoted. However, this effect cannot be obtained when the final cold rolling reduction is less than 40% or exceeds 80%, and is only achieved when the final cold rolling reduction is in the range of 40 to 80%.

After the cold rolling process as described above is completed, decarburization annealing is usually performed in a wet hydrogen atmosphere at a temperature range of 750 to 850°C to decarburize the carbon content to 0.003% or less. After that, an annealing separator such as MgO is applied, and then final annealing is performed. This final annealing is performed as follows:
In order to remove impurities such as S, Se, and N and to form an electrical insulating coating mainly made of forsterite,
It is desirable to hold the steel sheet at a temperature of about 900°C or higher, preferably in the range of 1050 to 1250°C, for several hours or more. When this final annealing is a high-temperature annealing of 900°C or higher, it is necessary to use a hydrogen atmosphere in order to promote the removal of impurities. However, when a low-temperature holding annealing at about 820 to 900°C is carried out before the high-temperature annealing, any of hydrogen, nitrogen, and argon may be used as the atmosphere.

Examples of the present invention are described below.

Example 1: A 200 mm thick continuous cast slab containing 3.15% Si, 0.045% C, 0.07% Mn, 0.025% S, and the balance essentially consisting of Fe, was heated to 1380°C for 1 hour, hot rolled to a thickness of _2.5 mm, and coiled. The hot rolled coil was then pickled and subjected to a first cold rolling to reduce the intermediate thickness to 0.70 mm. Subsequently, an intermediate annealing was performed at 925°C for 3 minutes to obtain a thickness of 0.70 _mm .
The decarburization amount ΔC was adjusted to three levels: 0.003%, which is less than the range of the present invention, 0.012%, which is within the range of the present invention, and 0.025%, which exceeds the range of the present invention. The subsequent cooling process was as follows: (A) the cooling time required between 770 and 300°C was 15 seconds, and then between 300 and 150°C for 15 seconds; (B) the cooling time required between 770 and 300°C was 15 seconds;
The steel sheets were cooled under two different conditions: a 60-second cooling time between 300 and 150°C, and a 15-second cooling time between 300 and 150°C. The steel sheets were then cold-rolled to a thickness of 0.30 mm with a reduction of 57%. The steel sheets were then decarburized in a wet hydrogen atmosphere at 800°C for 55 minutes, after which MgO slurry was applied and box annealed immediately at 1150°C.
The steel sheets were then heated to 1000°C, held for 15 hours for final annealing, and then coated with an insulating coating to obtain grain-oriented silicon steel sheets. The magnetic properties (magnetic flux density _B10 and iron loss _{W17 / 50} ) of these steel sheets were measured, and the results are shown in Table 3, corresponding to the process conditions.

In Table 3, both Samples 2 and 6 appear to have an appropriate amount of gamma phase formation within the range of 10-30% based on the amount of C in the raw material, but the amount of decarburization ΔC does not satisfy the range of 0.006-0.020% specified in this invention, and the carbide precipitate size is not within the range of 100-500 Å specified in this invention, so low iron loss values and high magnetic flux densities cannot be obtained.
Although the thickness of sample 1 was within the range of 0.00 Å, the amount of decarburization did not satisfy the range of this invention, so although the magnetic properties improved slightly, the desired sufficient characteristics were not obtained. On the other hand, sample 4 satisfied the amount of decarburization, but the carbide precipitate size did not satisfy the range of this invention, so similarly the magnetic properties improved slightly, but the desired sufficient characteristics were not obtained. In contrast, sample 3, which satisfied all the requirements of this invention, obtained a sufficiently low iron loss value and a high magnetic flux density at the same time.

Example 2 C 0.054%, Si 3.30%, Mn 0.85%, S 0.021%, Se
A 2mm thick hot-rolled sheet having a composition containing 0.010% of Cu and the balance essentially consisting of Fe was pickled and then annealed at 950°C for 2 minutes in a wet hydrogen atmosphere with P _{H2O /} P _H2 = 0.35 (amount of decarburization ΔC = 0.
After that, it was subjected to intermediate cold rolling to a thickness of 0.70 mm, and then further
The steel was then subjected to intermediate annealing at 770°C for 2 minutes (decarburization amount ΔC = 0.002%), and during the cooling process after the intermediate annealing, the steel was cooled between 770°C and 100°C in 15 seconds, and immediately aged at 200°C for 30 seconds.Then, the steel was finished to a thickness of 0.20 mm by final cold rolling with a reduction rate of 71%.
Thereafter, decarburization annealing was performed in a wet hydrogen atmosphere at 830°C for 3 minutes, and after applying MgO slurry, final annealing was performed, in which the material was held at 850°C for 50 hours during the temperature increase, and then purification annealing was performed at 1200°C for 10 hours. After that, an insulating coating was applied, and a grain-oriented silicon steel sheet product (Sample No. 8) of the present invention was obtained.

For comparison, a hot-rolled sheet having the same composition and thickness as the above-mentioned example of the present invention was annealed at 950°C for 2 minutes (decarburization amount ΔC = 0.00).
After that, it was pickled, cold-rolled to a thickness of 0.70 mm, and then annealed at 950°C for 2 minutes (amount of decarburization ΔC = 0.002%).
After the intermediate annealing, the cooling process is carried out at 770°C to 1
The cooling time between 0°C and 100°C was 50 seconds (no aging treatment was performed), and then the final cold rolling was performed at a rolling reduction of 71% to finish the product to a thickness of 0.20 mm. Thereafter, decarburization annealing, MgO slurry application, final annealing, and insulating coating were performed under the same conditions as in the above-mentioned example of the present invention, to produce a grain-oriented silicon steel sheet product for the comparative example (sample N).
o.7) was obtained.

The magnetic properties of these products were examined and the results are shown in Table 4.

As is clear from Table 4, the product of the present invention (sample No. 8) manufactured by the method of the present invention was found to have superior magnetic properties compared to the product of the comparative example (sample No. 7).

As is clear from the above explanation, according to the manufacturing method of the present invention, the C content of the material is adjusted to an appropriate range depending on the Si content, the amount of decarburization from hot rolling to before final cold rolling is kept within an appropriate range, and the carbides within the crystal grains of the steel sheet before final cold rolling are appropriately controlled, thereby achieving a significantly high magnetic flux density, which has not been achieved conventionally.
It is possible to stably obtain grain-oriented silicon steel sheets having extremely excellent magnetic properties with extremely low iron loss values, and also in terms of the process, grain-oriented silicon steel sheets with extremely excellent properties can be obtained without requiring special slow cooling at high temperatures or long-term aging treatment, so that various effects can be obtained, such as high productivity and economical efficiency even when implemented on an industrial scale.

[Brief explanation of the drawings]

FIG. 1 is a graph showing the influence of the amount of Si and the amount of C contained in the material on the iron loss value of the product, FIG. 2 is a graph showing the influence of the amount of decarburization ΔC after hot rolling and before final cold rolling on the magnetic properties of the product, FIG. 3 is a graph showing the relationship between the amount of decarburization in intermediate annealing and the aging treatment time when quenched after intermediate annealing and aging treatment at 200°C, and the magnetic properties and carbide precipitation size, and FIG. 4 is a metallographic photograph taken by an electron microscope at a magnification of 10,000 times to show the carbide precipitation state of the steel sheet before final cold rolling.
(A) shows the results when quenching and aging treatment are performed after intermediate annealing according to the present invention, and (B) shows the results when standard cooling after intermediate annealing is performed according to the conventional method. FIG. 5 is a graph comparing the Goss orientation strength of the surface layer of the steel sheet after decarburization annealing depending on whether or not decarburization is performed in the intermediate annealing step and whether or not quenching-aging treatment is performed after intermediate annealing. FIG. 6 is a graph showing the relationship between the aging treatment conditions and the size of carbide precipitates when quenching is performed after intermediate annealing and then aging treatment is performed before final cold rolling.
The figure shows the results for the cooling process after intermediate annealing before final cold rolling.
1 is a graph showing the relationship between the cooling time required between 300 and 150°C, the carbide precipitation size, and the magnetic properties when rapid cooling is performed between 300 and 150°C and the cooling time required between 300 and 150°C is changed.

Claims (2)

[Claims] Claim 1: C 0.015 to 0.10% (weight %, the same applies hereinafter),
Contains 2.8 to 4.0% Si, 0.02 to 0.15% Mn, and S, S
One or two of the above e in total amount of 0.008 to 0.080%
In a series of manufacturing methods for grain-oriented silicon steel sheet, a silicon steel material containing 0.37[Si%]+0.27≦log([C%]×103)≦0.37 is hot-rolled, the resultant hot-rolled steel sheet is cold-rolled at least twice with intermediate annealing in between at a final cold-rolling reduction of 40-80% to be finished to a predetermined thickness, and the cold-rolled sheet is then decarburized and final annealed, and the amount of C contained in the silicon steel material is determined in accordance with the amount of Si in accordance with the following formula: 0.37[Si%]+0.27≦log([C%]× ¹⁰³ )≦0.37
[Si%] + 0.57, and after hot rolling is completed,
A method for producing grain oriented silicon steel sheet with excellent magnetic properties, characterized by decarburizing 0.006 to 0.020% of C before the end of final cold rolling, and rapidly cooling the steel sheet after intermediate annealing before final cold rolling so that the cooling time between 770 and 100°C is within 30 seconds, followed by aging treatment for 2 to 60 seconds in a temperature range of 150 to 250°C, and then final cold rolling. Claim 2: C 0.015 to 0.10% (weight %, the same applies hereinafter),
Contains 2.8 to 4.0% Si, 0.02 to 0.15% Mn, and S, S
One or two of the above e in total amount of 0.008 to 0.080%
In a series of manufacturing methods for grain-oriented silicon steel sheet, a silicon steel material containing 0.37[Si%]+0.27≦log([C%]×103)≦0.37 is hot-rolled, the resultant hot-rolled steel sheet is cold-rolled at least twice with intermediate annealing in between at a final cold-rolling reduction of 40-80% to be finished to a predetermined thickness, and the cold-rolled sheet is then decarburized and final annealed, and the amount of C contained in the silicon steel material is determined in accordance with the amount of Si in accordance with the following formula: 0.37[Si%]+0.27≦log([C%]× ¹⁰³ )≦0.37
[Si%] + 0.57, and after hot rolling is completed,
A method for producing grain oriented silicon steel sheet with excellent magnetic properties, characterized by decarburizing the carbon by 0.006 to 0.020% before the end of final cold rolling, and controlling the cooling time required between 770 and 300°C within 20 seconds in the cooling process after intermediate annealing before final cold rolling, and subsequently controlling the cooling time required between 300 and 150°C within a range of 8 to 30 seconds, before final cold rolling.