CA2048014C - Method of manufacturing an oriented silicon steel sheet having improved magnetic flux density - Google Patents
Method of manufacturing an oriented silicon steel sheet having improved magnetic flux densityInfo
- Publication number
- CA2048014C CA2048014C CA 2048014 CA2048014A CA2048014C CA 2048014 C CA2048014 C CA 2048014C CA 2048014 CA2048014 CA 2048014 CA 2048014 A CA2048014 A CA 2048014A CA 2048014 C CA2048014 C CA 2048014C
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- Prior art keywords
- steel sheet
- temperature
- sheet
- annealing
- rolling
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Classifications
-
- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
- C21D8/12—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of articles with special electromagnetic properties
- C21D8/1244—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of articles with special electromagnetic properties the heat treatment(s) being of interest
- C21D8/125—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of articles with special electromagnetic properties the heat treatment(s) being of interest with application of tension
-
- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
- C21D8/12—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of articles with special electromagnetic properties
- C21D8/1216—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of articles with special electromagnetic properties the working step(s) being of interest
- C21D8/1227—Warm rolling
-
- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
- C21D8/12—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of articles with special electromagnetic properties
- C21D8/1244—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of articles with special electromagnetic properties the heat treatment(s) being of interest
- C21D8/1261—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of articles with special electromagnetic properties the heat treatment(s) being of interest following hot rolling
-
- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
- C21D8/12—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of articles with special electromagnetic properties
- C21D8/1244—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of articles with special electromagnetic properties the heat treatment(s) being of interest
- C21D8/1266—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of articles with special electromagnetic properties the heat treatment(s) being of interest between cold rolling steps
Abstract
A method of manufacturing an oriented silicon steel sheet which achieves a high magnetic flux density while reducing the core loss. A silicon steel sheet containing Al and Sb as inhibitor components is cold-rolled once or a plurality of times. During cooling for annealing before final cold rolling, a small strain is created on the sheet and the temperature is within a certain range. Carbide precipitation is suitably controlled to precipitate carbides comparatively coarsely in grains.
Description
20~01~
R~GROUND OF THE lNv~NlION
This invention relates to a method of manufacturing an oriented silicon steel sheet having improved magnetic characteristics and, more particularly, to an improved method of preventing reduction of magnetic flux density notwithstanding reduction of thickness of the silicon steel sheet.
High magnetic flux density and a small core loss are magnetic characteristics required in grain-oriented silicon steel sheets. Recent progress in manufacture techniques has made it possible to make, for example, a silicon steel sheet having a magnetic flux density B8 (the value at a magnetizing force of 800 A/m) of 1.92 T for a sheet having a thickness of 0.23 mm. It is also possible to manufacture, on an industrial scale, an improved silicon steel sheet product having a core loss characteristic W"/50 (value under a fully magnetized condition: 1.7 T at 50 Hz) of 0.90 w/kg.
Silicon steel sheets having such improved magnetic characteristics have crystalline structures in which the <001>
directions parallel to the axis of easy magnetization are uniformly aligned in the direction of rolling of the steel sheet. Such a texture is formed during finishing annealing by a phenomenon called secondary recrystallization in which crystal grains having a (110) ~001] direction called the Goss directian are grown with priority into giant grains.
Fundamental requirements for effectively growing secondary recrystallized grains include the existence of an inhibitor for limiting the growth of crystal grains having undesirable directions other than the (110) [001] direction in the secondary recrystallization process and the formation of a primary recrystallized crystalline structure suitable for effectively developing secondary recrystallized grains in the (110) [001] direction.
A fine precipitate of MnS, MnSe, AlN or the like is ordinarily utilized as the inhibitor. The effect of the inhibitor has been enhanced by adding a grain boundary segregation type component such as Sb or Sn to the inhibitor.
Conventionally, methods in which NnS or MnSe is used as a main inhibitor are advantageous in reducing the core loss of certain sheets because they assist in reducing the sizes of the secondary recrystallized grains. However, methods based on laser irradiation or plasma jetting have recently been provided to artificially form pseudo grain boundaries so that the magnetic domains are fractionated and the core loss is reduced. For this reason, the advantage of reducing the sizes of the secondary recrystallized grains has been lost.
Further, the concept of increasing the magnetic flux density of the steel sheet has become advantageous.
A method of manufacturing an oriented silicon steel sheet having a large magnetic flux density is disclosed in Japanese patent Publication 46-23820. According to this method, the desired steel sheet can be manufactured by (a) introducing Al into the steel as an inhibitor component, (b) quenching to obtain cooling before final cold rolling to precipitate AlN, and (c) increasing the rolling reduction of the final cold rolling from a lower reduction to a higher reduction, like from 65 to 95 %.
The method of the Japanese Publication, however, entails a problem in that the magnetic flux is abruptly reduced along with the reduction of thickness of the product sheet. It is very difficult or impossible to manufacture by the method of the Japanese Publication the type of silicon steel sheet presently in demand, e.g., a thin product having a thickness of 0.25 mm or less and having a B8 value of 1.94 T or higher.
In Japanese patent Publication 46-23820, immersing a steel sheet in hot water at 100C after annealing to quench the sheet is disclosed, but there is no consideration or mention of any phase of any carbides after quenching.
Ordinarily, in the case of slow cooling from 600C or lower, carbides are precipitated from grain boundaries at a higher temperature and are precipitated in crystal grains at a lower temperature. Carbides precipitated are finer and have a higher density if precipitation is started at a reduced temperature. Accordingly, with respect to the first embodiment of Japanese patent Publication 46-23820 in which the time for cooling from 1,000 to 750C is about 10 seconds and the time for cooling from 750 to 100C is about 25 seconds, it is not unreasonable to conclude that very fine carbides having particle sizes of several tens of angstroms are precipitated or that the extent of carbide precipitation is limited and that the carbon is simply supersaturated in the steel.
Japanese Patent Publication 56-3892 discloses a technique for controlling carbides in other steels during cooling after 204801~
annealing. In this method, with respect to two-stage cold rolling, the steel is cooled at a cooling speed of 150C/min or higher from 600 to 300C during cooling after annealing followed by final cold rolling so that the amount of solid solution carbon after cooling is increased. This method is intended to improve the magnetic characteristics of the steel by increasing the amount of solid solution carbon in the steel and by optimizing the aging effect between cold rolling paths.
Such an effect of solid solution carbon is well known in the case of ordinary cold-rolled steel sheets. If the amount of solid solution C or solid solution N before cold rolling is increased, the (110) intensity in the recrystallized structure formed by recrystallization annealing after cold rolling is increased. In the case of oriented silicon steel sheets, the (110) grains become nuclei for secondary recrystallization, so that the number of secondary recrystallized grains is increased, the secondary-recrystallized grains are finer, and improved magnetic characteristics can be achieved. This method, however, does not enable the magnetic flux density of a thin oriented silicon steel sheet to be increased.
As a technique for controlling the form of C in steel to increase the (110) intensity of the steel, a method of precipitating many fine carbide grains during cooling after intermediate annealing is disclosed in Japanese Patent Laid-Open Publication 58-157917. In this method, quenching of the steel to 300C is effected after intermediate annealing and slow cooling is applied for 8 to 30 seconds through a temperature range of 300 to 150C, thereby precipitating fine carbides. The (110) intensity of the steel after recrystallization is thereby increased so that the magnetic characteristics of the steel are improved. However, the magnetic characteristics achieved by these methods are at most 1.94 T with respect to Blo and 1.92 T with respect to B8 when the sheet thickness is 0.3 mm, which value is not high enough to be satisfactory.
Japanese Patent Laid-Open Publication 61-149432discloses a technique based on setting the cooling speed of steel to 10C/s or higher at the time of cooling after intermediate annealing, creating a work strain of 1 to 30 % during cooling from 1,000 to 400C, and perorming finishing rolling at a temperature in the range of 100C to 400C. According to this method, a work strain of 1 to 30 % is created at a temperature in the range of 1,000 to 400C in which the C diffusion speed is very high to provide high-density dislocations, so that C
is finely precipitated at the dislocations and the (110) intensity is increased. To finely precipitate C in dislocations at a high density, the working is performed by rolling, and a high cooling speed of 10C/s or higher is set for the precipitation step. The core loss can be reduced to a certain extent by this method but the magnetic flux density achieved by this method is only 1.91 T with respect to B1o (1.89 T with respect to B8), which is low.
20~80~4 OBJECTS OF THE INVENTION
It is an object of the present invention to provide a method of manufacturing an oriented silicon steel sheet which enables maintenance of high magnetic flux density notwithstanding reduction of steel sheet thickness. Another object is to achieve a high magnetic flux density with desired stability while reducing the core loss of steel sheet.
SUMM~RY OF THE lNV~;N'l'ION
It has been discovered that, in an Al-containing oriented silicon steel sheet in which Sb is also present, the precipitation of carbides is greatly changed during cooling for annealing before final cold rolling, and that such precipitation is effective to increase the ultimate (lll) intensity of the recrystallized structure after final cold rolling of sheet rather than the (110) intensity, and that carbides precipitated in crystal grains at a high temperature in the range of about 200 to 500C under strain during cooling for annealing before final cold rolling, which are conventionally regarded as undesirable, surprisingly have the effect of increasing the {111} <112> intensity while reducing the {111} <uvw> intensity, more particularly the {111} <llO>
intensity, so that a very high magnetic flux density can be obtained with stability irrespective of the thickness of the final product.
That is, according to the present invention, there is provided a method of manufacturing an oriented silicon steel sheet having greatly improved magnetic characteristics in which a hot-rolled steel sheet of a silicon steel containing about 0.01 to 0.15 6 by wel~ht of acid-solub.Le Al and abo~lt 0.005 to 0.04 % by weight of Sb as lnhlbltor components ls cold-rolled once or a plurality of times until its thickness ls reduced to the desired predetermined final thickness. The method further comprises softening-annealing the steel sheet before final cold-rolling;
successlvely guenchlng the steel sheet at a cooling speed of about 15 to 500C/s to a temperature of about 500C
or lower;
creatlng upon the sheet a small strain ranglng from about 0.005 to 3.0 % ln a temperature range from about the temperature reached by quenching to about 200C;
controlling carblde preclpltation at an effective coollng speed about 2~/s or less by cooling the steel sheet durlng the stralnlng step or by coollng the steel sheet for about 60 to 180 seconds after the strainlng step durinq whlch the steel sheet ls malntalned ln essentlally the same temperature range as the temperature range at the end of the quenchlng step; and thereafter performing final cold-rolling with a rolllng reductlon of about 80 to 95 %.
Thls can be done in coniunctlon with additional steps of effecting anneallng for primary recrystalllzatlon as well as decarburization;
applying an anneallng separation agent; and effectlng secondary recrystalllzatlon annealing and purlflcatlon-annealing.
73461~27 E ~!
Other features and variations of the present invention will become apparent from the following detailed description of the invention.
BRIEF DESCRIPTION OF THE DRAWINGS
Figs. 1 to 4 are transmission-electron-microscopic photographs of examples of structures of steel sheets after 8a t ~ 73461-27 ii '.V
annealing followed by final cold rolling, showing forms of carbides at a depth of one-tenth of the sheet thickness measured from the surfaces of the steel sheets.
DETAILED DESCRIPTION OF THE INVENTION
First, the results of experiments on which the present invention is based will be described below.
Al-containing oriented silicon steel sheets to which Sb, Sn, Ge, Ni and Cu (well-known as additive components) were separately added were provided. These sheets were rolled different times to manufacture products; one group of these steel sheets was cold-rolled only one time to obtain products having a thickness of 0.30 mm, and another group was cold-rolled twice to obtain products having a thickness of 0.23 mm.
The rolling reduction of the final cold rolling was set at 88 %, and annealing immediately before final cold rolling was performed at 1,150C for 90 seconds with respect to the steel sheets cold-rolled one time, and at 1,100C for 90 seconds with respect to the steel sheets cold-rolled twice.
Cooling was performed by immersing each steel sheet in hot water at 80C.
The results of this experiment are as shown in Table 1.
Each of the 0.30 mm thick steel sheets had a high magnetic flux density while each of the 0.23 mm thick steel sheets had a reduced magnetic flux density. The reduced sheet thickness had seriously reduced the flux density in every case.
Table 1 Sample No. 1 2 3 4 5 6 Additive Constituent name No additive Ni Cu Sb Sn Ge - Amount of additive (%) - 0.08 0.10 0.03 0.05 0.02 Magnetic flux Product 0.30mm 1.924 1.925 1.923 1.936 1.903 1.914 density B8 (T) thickness 0.23mm 1.885 1.885 1.887 1.894 1.882 1.884 O O
o 204801~
By ex~mi ning the results of Table 1 in detail, it is evident that sample 4 in which Sb was present had a slightly better magnetic flux density than the other five samples.
To ~x~mine the cause of this effect we examined the textures of samples of decarburized primary recrystallized sheets with respect to the samples having a product thickness of 0.23 mm, and ex~m; ~ed the forms of precipitated carbides in the steel of each sample after intermediate annealing. The results of these ex~mi n~tions are shown in Table 2.
Table 2 Sample No. 1 2 3 4 5 6 Additive No additive Ni Cu Sb Sn Ge constituent (110) Intensity 0.15 0.16 0.18 0.12 0.22 0.25 (222) Intensity 7.3 7.5 7.0 8.8 6.4 6.8 Form of carbideMostly in solid Mostly in solid PrecipitatedPrecipitated finely O
precipitation in solution, partially solution, partially slightly coarselyand at a high ~, intermediate precipitated finely precipitated finely in grainsdensity in grains annealed sheet ~, Precipitated sizeAbout 80~ About 80~ About 200~ About 50~
As can be understood from Table 2, no increase in the (110) intensity is attributed to the presence of Sb as observed in sample 4 containing Sb, unlike the effect that might have been expected in view of conventional technical concepts, but the (111) intensity (equivalent to (222)) was remarkably increased in the sample containing Sb. Further, different forms of carbides exist after annealing followed by final cold rolling and, as a result of the addition of Sb, the fine high-density precipitated state or the C solid solution state was changed so that carbides were precipitated in the form of slightly coarse grains (Table 2, column 4) having particle sizes much greater than the others in the Table.
In contrast, in the case of addition of Sn or Ge, carbides were finely precipitated at a high density, and the (110) intensity of the primary recrystallized structure was remarkably improved.
The cause of this special effect achieved by the presence of Sb is not clear. However, it is speculated that the tendency of Sb to strongly segregate at grain boundaries or surfaces is related to the phenomenon leading to the occurrence of specially precipitated forms of carbides.
With a view to positive utilization of such variations of the forms of carbides before final cold rolling, and to create new effects by varying cooling conditions, further experiments were conducted. Tests were conducted on the same Al-containing oriented silicon steel sheets as those used in the above-described experiments to which only Sb was added, and also on the same Al-contAining silicon steel sheets which had no added component. The tested steels were processed by ordinary two-stage rolling to product products each having a thickness of 0.23 mm. In this experiment, the rolling reduction of final cold rolling was set at 85 ~, annealing before the final cold rolling (intermediate annealing) was effected at 1,100C for 90 seconds, and cooling was effected under the following different cooling conditions:
(a) Condition (a) wherein the steel sheet was quenched at a rate of 50C/s until 500C was reached, and thereafter cooled at a very low cooling speed of 0.5 to 2C/s by being inserted in a heat maintaining furnace, (b) Condition (b) wherein the steel sheet was quenched at a rate of 50C/s until 350C was reached, and thereafter cooled at a very low cooling speed of 0.5 to 2C/s by insertion into a heat maintaining furnace, (c) Condition (c) wherein the steel sheet was quenched at a rate of 50C/s until 350C was reached, successively skin-pass-rolled to reduce by O.S ~, and cooled at a very low cooling speed of 0.5 to 2C/s by insertion into a heat maintAi~in~ furnace, (d) Condition (d) wherein the steel sheet was quenched at a rate of 50C/s until 150C was reached, and thereafter cooled at a very low cooling speed of 0.5 to 2C/s by insertion into a heat maintaining furnace, (e) Condition (e) wherein the steel sheet was immersed in hot water at 80C so that the average cooling speed was 62C/s, was maintained at 80C after being cooled to this temperature, and was thereafter cooled naturally.
The products thereby manufactured were examined with respect to magnetic flux density, (110) intensity and (222) intensity of the decarburized primary recrystallized sheets and the precipitated forms of carbides in the intermediate annealed sheets. The results are shown in Table 3.
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i C' ~ ci~ c' ~: o ~ o ~ u Figs. 1 to 4 are transmission-electron-microscopic photographs of the structures of steel sheets after annealing followed by final cold rolling, showing forms of carbides at a depth of 1/10 of the sheet thickness from the surfaces of the steel sheets. Fig. 1 shows a sample to which Sb was added and which was cooled under Condition (e), Fig. 2 shows a sample (Table 3, column c, bottom) to which Sb was added and which was cooled under Condition (c), Fig. 3 shows a sample which had no additive component and which was cooled under Condition (e) (Table 3, column 3, top), and Fig. 4 shows a sample which had no additive component and which was cooled under Condition (c) (Table 3, column c, top).
As is shown in Table 3, the magnetic flux density (B8)(T) of the sample to which Sb was added (bottom half of Table 3) and which was manufactured under the intermediate annealing cooling condition (c) (Table 3, column c) was particularly high. In this sample, carbide precipitates having a size ranging from 300 to 500 A and sparsely precipitated were observed after the intermediate annealing, and are shown in Fig. 2, as heretofore noted. In contrast, in the sample which had no additive component and which was manufactured under the same cooling condition (c) (Table 3, column c, top), fine carbide precipitates having a size of about 100 A were undesirably precipitated at a high density, as shown in ~ig.
4.
With respect to the steel sheets which had no additive - component, in the case of creating a work strain by skin-pass-rolling in accordance with the condition (c), carbide precipitation sites were increased during cooling so that carbides were finely precipitated at a high density, as is apparent from comparison with processing under the condition (b). In contrast, with respect to the steel sheets to which Sb was added, precipitation sites were not increased and slightly coarse precipitates were observed. According to our study after these experiments, such sparse precipitation of carbides having a size ranging from 300 to 500 ~ increases the (111) intensity of the structure primarily recrystallized by decarburization annealing after final cold rolling and reduces the {lll} <uvw>, in particular the {111} <110> intensity while increasing the {111} <112> intensity. The {111} <110> grains limit the growth of the (110) [001] secondary grains which contribute to the increase in the magnetic flux density, while the {111} <112> grains promote the growth of (110) [001]
secondary grains. It is thought that addition of Sb in the particular process provides this effect and enables formation of a product having a substantially high magnetic flux density as in the case of Condition (c) as shown in the top portion of Table 3.
It is thought that this effect of Sb in steel relates to segregation of Sb, that Sb is segregated at base points in crystal grains such as to form carbide precipitation sites, and that this segregation results from the limitation of precipitation carbides during cooling.
This action of Sb is particularly effective in a temperature range of about 200 to 500C; the amount of strain to be applied may be very small, e.g., about 0.005 to 3 %.
R~GROUND OF THE lNv~NlION
This invention relates to a method of manufacturing an oriented silicon steel sheet having improved magnetic characteristics and, more particularly, to an improved method of preventing reduction of magnetic flux density notwithstanding reduction of thickness of the silicon steel sheet.
High magnetic flux density and a small core loss are magnetic characteristics required in grain-oriented silicon steel sheets. Recent progress in manufacture techniques has made it possible to make, for example, a silicon steel sheet having a magnetic flux density B8 (the value at a magnetizing force of 800 A/m) of 1.92 T for a sheet having a thickness of 0.23 mm. It is also possible to manufacture, on an industrial scale, an improved silicon steel sheet product having a core loss characteristic W"/50 (value under a fully magnetized condition: 1.7 T at 50 Hz) of 0.90 w/kg.
Silicon steel sheets having such improved magnetic characteristics have crystalline structures in which the <001>
directions parallel to the axis of easy magnetization are uniformly aligned in the direction of rolling of the steel sheet. Such a texture is formed during finishing annealing by a phenomenon called secondary recrystallization in which crystal grains having a (110) ~001] direction called the Goss directian are grown with priority into giant grains.
Fundamental requirements for effectively growing secondary recrystallized grains include the existence of an inhibitor for limiting the growth of crystal grains having undesirable directions other than the (110) [001] direction in the secondary recrystallization process and the formation of a primary recrystallized crystalline structure suitable for effectively developing secondary recrystallized grains in the (110) [001] direction.
A fine precipitate of MnS, MnSe, AlN or the like is ordinarily utilized as the inhibitor. The effect of the inhibitor has been enhanced by adding a grain boundary segregation type component such as Sb or Sn to the inhibitor.
Conventionally, methods in which NnS or MnSe is used as a main inhibitor are advantageous in reducing the core loss of certain sheets because they assist in reducing the sizes of the secondary recrystallized grains. However, methods based on laser irradiation or plasma jetting have recently been provided to artificially form pseudo grain boundaries so that the magnetic domains are fractionated and the core loss is reduced. For this reason, the advantage of reducing the sizes of the secondary recrystallized grains has been lost.
Further, the concept of increasing the magnetic flux density of the steel sheet has become advantageous.
A method of manufacturing an oriented silicon steel sheet having a large magnetic flux density is disclosed in Japanese patent Publication 46-23820. According to this method, the desired steel sheet can be manufactured by (a) introducing Al into the steel as an inhibitor component, (b) quenching to obtain cooling before final cold rolling to precipitate AlN, and (c) increasing the rolling reduction of the final cold rolling from a lower reduction to a higher reduction, like from 65 to 95 %.
The method of the Japanese Publication, however, entails a problem in that the magnetic flux is abruptly reduced along with the reduction of thickness of the product sheet. It is very difficult or impossible to manufacture by the method of the Japanese Publication the type of silicon steel sheet presently in demand, e.g., a thin product having a thickness of 0.25 mm or less and having a B8 value of 1.94 T or higher.
In Japanese patent Publication 46-23820, immersing a steel sheet in hot water at 100C after annealing to quench the sheet is disclosed, but there is no consideration or mention of any phase of any carbides after quenching.
Ordinarily, in the case of slow cooling from 600C or lower, carbides are precipitated from grain boundaries at a higher temperature and are precipitated in crystal grains at a lower temperature. Carbides precipitated are finer and have a higher density if precipitation is started at a reduced temperature. Accordingly, with respect to the first embodiment of Japanese patent Publication 46-23820 in which the time for cooling from 1,000 to 750C is about 10 seconds and the time for cooling from 750 to 100C is about 25 seconds, it is not unreasonable to conclude that very fine carbides having particle sizes of several tens of angstroms are precipitated or that the extent of carbide precipitation is limited and that the carbon is simply supersaturated in the steel.
Japanese Patent Publication 56-3892 discloses a technique for controlling carbides in other steels during cooling after 204801~
annealing. In this method, with respect to two-stage cold rolling, the steel is cooled at a cooling speed of 150C/min or higher from 600 to 300C during cooling after annealing followed by final cold rolling so that the amount of solid solution carbon after cooling is increased. This method is intended to improve the magnetic characteristics of the steel by increasing the amount of solid solution carbon in the steel and by optimizing the aging effect between cold rolling paths.
Such an effect of solid solution carbon is well known in the case of ordinary cold-rolled steel sheets. If the amount of solid solution C or solid solution N before cold rolling is increased, the (110) intensity in the recrystallized structure formed by recrystallization annealing after cold rolling is increased. In the case of oriented silicon steel sheets, the (110) grains become nuclei for secondary recrystallization, so that the number of secondary recrystallized grains is increased, the secondary-recrystallized grains are finer, and improved magnetic characteristics can be achieved. This method, however, does not enable the magnetic flux density of a thin oriented silicon steel sheet to be increased.
As a technique for controlling the form of C in steel to increase the (110) intensity of the steel, a method of precipitating many fine carbide grains during cooling after intermediate annealing is disclosed in Japanese Patent Laid-Open Publication 58-157917. In this method, quenching of the steel to 300C is effected after intermediate annealing and slow cooling is applied for 8 to 30 seconds through a temperature range of 300 to 150C, thereby precipitating fine carbides. The (110) intensity of the steel after recrystallization is thereby increased so that the magnetic characteristics of the steel are improved. However, the magnetic characteristics achieved by these methods are at most 1.94 T with respect to Blo and 1.92 T with respect to B8 when the sheet thickness is 0.3 mm, which value is not high enough to be satisfactory.
Japanese Patent Laid-Open Publication 61-149432discloses a technique based on setting the cooling speed of steel to 10C/s or higher at the time of cooling after intermediate annealing, creating a work strain of 1 to 30 % during cooling from 1,000 to 400C, and perorming finishing rolling at a temperature in the range of 100C to 400C. According to this method, a work strain of 1 to 30 % is created at a temperature in the range of 1,000 to 400C in which the C diffusion speed is very high to provide high-density dislocations, so that C
is finely precipitated at the dislocations and the (110) intensity is increased. To finely precipitate C in dislocations at a high density, the working is performed by rolling, and a high cooling speed of 10C/s or higher is set for the precipitation step. The core loss can be reduced to a certain extent by this method but the magnetic flux density achieved by this method is only 1.91 T with respect to B1o (1.89 T with respect to B8), which is low.
20~80~4 OBJECTS OF THE INVENTION
It is an object of the present invention to provide a method of manufacturing an oriented silicon steel sheet which enables maintenance of high magnetic flux density notwithstanding reduction of steel sheet thickness. Another object is to achieve a high magnetic flux density with desired stability while reducing the core loss of steel sheet.
SUMM~RY OF THE lNV~;N'l'ION
It has been discovered that, in an Al-containing oriented silicon steel sheet in which Sb is also present, the precipitation of carbides is greatly changed during cooling for annealing before final cold rolling, and that such precipitation is effective to increase the ultimate (lll) intensity of the recrystallized structure after final cold rolling of sheet rather than the (110) intensity, and that carbides precipitated in crystal grains at a high temperature in the range of about 200 to 500C under strain during cooling for annealing before final cold rolling, which are conventionally regarded as undesirable, surprisingly have the effect of increasing the {111} <112> intensity while reducing the {111} <uvw> intensity, more particularly the {111} <llO>
intensity, so that a very high magnetic flux density can be obtained with stability irrespective of the thickness of the final product.
That is, according to the present invention, there is provided a method of manufacturing an oriented silicon steel sheet having greatly improved magnetic characteristics in which a hot-rolled steel sheet of a silicon steel containing about 0.01 to 0.15 6 by wel~ht of acid-solub.Le Al and abo~lt 0.005 to 0.04 % by weight of Sb as lnhlbltor components ls cold-rolled once or a plurality of times until its thickness ls reduced to the desired predetermined final thickness. The method further comprises softening-annealing the steel sheet before final cold-rolling;
successlvely guenchlng the steel sheet at a cooling speed of about 15 to 500C/s to a temperature of about 500C
or lower;
creatlng upon the sheet a small strain ranglng from about 0.005 to 3.0 % ln a temperature range from about the temperature reached by quenching to about 200C;
controlling carblde preclpltation at an effective coollng speed about 2~/s or less by cooling the steel sheet durlng the stralnlng step or by coollng the steel sheet for about 60 to 180 seconds after the strainlng step durinq whlch the steel sheet ls malntalned ln essentlally the same temperature range as the temperature range at the end of the quenchlng step; and thereafter performing final cold-rolling with a rolllng reductlon of about 80 to 95 %.
Thls can be done in coniunctlon with additional steps of effecting anneallng for primary recrystalllzatlon as well as decarburization;
applying an anneallng separation agent; and effectlng secondary recrystalllzatlon annealing and purlflcatlon-annealing.
73461~27 E ~!
Other features and variations of the present invention will become apparent from the following detailed description of the invention.
BRIEF DESCRIPTION OF THE DRAWINGS
Figs. 1 to 4 are transmission-electron-microscopic photographs of examples of structures of steel sheets after 8a t ~ 73461-27 ii '.V
annealing followed by final cold rolling, showing forms of carbides at a depth of one-tenth of the sheet thickness measured from the surfaces of the steel sheets.
DETAILED DESCRIPTION OF THE INVENTION
First, the results of experiments on which the present invention is based will be described below.
Al-containing oriented silicon steel sheets to which Sb, Sn, Ge, Ni and Cu (well-known as additive components) were separately added were provided. These sheets were rolled different times to manufacture products; one group of these steel sheets was cold-rolled only one time to obtain products having a thickness of 0.30 mm, and another group was cold-rolled twice to obtain products having a thickness of 0.23 mm.
The rolling reduction of the final cold rolling was set at 88 %, and annealing immediately before final cold rolling was performed at 1,150C for 90 seconds with respect to the steel sheets cold-rolled one time, and at 1,100C for 90 seconds with respect to the steel sheets cold-rolled twice.
Cooling was performed by immersing each steel sheet in hot water at 80C.
The results of this experiment are as shown in Table 1.
Each of the 0.30 mm thick steel sheets had a high magnetic flux density while each of the 0.23 mm thick steel sheets had a reduced magnetic flux density. The reduced sheet thickness had seriously reduced the flux density in every case.
Table 1 Sample No. 1 2 3 4 5 6 Additive Constituent name No additive Ni Cu Sb Sn Ge - Amount of additive (%) - 0.08 0.10 0.03 0.05 0.02 Magnetic flux Product 0.30mm 1.924 1.925 1.923 1.936 1.903 1.914 density B8 (T) thickness 0.23mm 1.885 1.885 1.887 1.894 1.882 1.884 O O
o 204801~
By ex~mi ning the results of Table 1 in detail, it is evident that sample 4 in which Sb was present had a slightly better magnetic flux density than the other five samples.
To ~x~mine the cause of this effect we examined the textures of samples of decarburized primary recrystallized sheets with respect to the samples having a product thickness of 0.23 mm, and ex~m; ~ed the forms of precipitated carbides in the steel of each sample after intermediate annealing. The results of these ex~mi n~tions are shown in Table 2.
Table 2 Sample No. 1 2 3 4 5 6 Additive No additive Ni Cu Sb Sn Ge constituent (110) Intensity 0.15 0.16 0.18 0.12 0.22 0.25 (222) Intensity 7.3 7.5 7.0 8.8 6.4 6.8 Form of carbideMostly in solid Mostly in solid PrecipitatedPrecipitated finely O
precipitation in solution, partially solution, partially slightly coarselyand at a high ~, intermediate precipitated finely precipitated finely in grainsdensity in grains annealed sheet ~, Precipitated sizeAbout 80~ About 80~ About 200~ About 50~
As can be understood from Table 2, no increase in the (110) intensity is attributed to the presence of Sb as observed in sample 4 containing Sb, unlike the effect that might have been expected in view of conventional technical concepts, but the (111) intensity (equivalent to (222)) was remarkably increased in the sample containing Sb. Further, different forms of carbides exist after annealing followed by final cold rolling and, as a result of the addition of Sb, the fine high-density precipitated state or the C solid solution state was changed so that carbides were precipitated in the form of slightly coarse grains (Table 2, column 4) having particle sizes much greater than the others in the Table.
In contrast, in the case of addition of Sn or Ge, carbides were finely precipitated at a high density, and the (110) intensity of the primary recrystallized structure was remarkably improved.
The cause of this special effect achieved by the presence of Sb is not clear. However, it is speculated that the tendency of Sb to strongly segregate at grain boundaries or surfaces is related to the phenomenon leading to the occurrence of specially precipitated forms of carbides.
With a view to positive utilization of such variations of the forms of carbides before final cold rolling, and to create new effects by varying cooling conditions, further experiments were conducted. Tests were conducted on the same Al-containing oriented silicon steel sheets as those used in the above-described experiments to which only Sb was added, and also on the same Al-contAining silicon steel sheets which had no added component. The tested steels were processed by ordinary two-stage rolling to product products each having a thickness of 0.23 mm. In this experiment, the rolling reduction of final cold rolling was set at 85 ~, annealing before the final cold rolling (intermediate annealing) was effected at 1,100C for 90 seconds, and cooling was effected under the following different cooling conditions:
(a) Condition (a) wherein the steel sheet was quenched at a rate of 50C/s until 500C was reached, and thereafter cooled at a very low cooling speed of 0.5 to 2C/s by being inserted in a heat maintaining furnace, (b) Condition (b) wherein the steel sheet was quenched at a rate of 50C/s until 350C was reached, and thereafter cooled at a very low cooling speed of 0.5 to 2C/s by insertion into a heat maintaining furnace, (c) Condition (c) wherein the steel sheet was quenched at a rate of 50C/s until 350C was reached, successively skin-pass-rolled to reduce by O.S ~, and cooled at a very low cooling speed of 0.5 to 2C/s by insertion into a heat maintAi~in~ furnace, (d) Condition (d) wherein the steel sheet was quenched at a rate of 50C/s until 150C was reached, and thereafter cooled at a very low cooling speed of 0.5 to 2C/s by insertion into a heat maintaining furnace, (e) Condition (e) wherein the steel sheet was immersed in hot water at 80C so that the average cooling speed was 62C/s, was maintained at 80C after being cooled to this temperature, and was thereafter cooled naturally.
The products thereby manufactured were examined with respect to magnetic flux density, (110) intensity and (222) intensity of the decarburized primary recrystallized sheets and the precipitated forms of carbides in the intermediate annealed sheets. The results are shown in Table 3.
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i C' ~ ci~ c' ~: o ~ o ~ u Figs. 1 to 4 are transmission-electron-microscopic photographs of the structures of steel sheets after annealing followed by final cold rolling, showing forms of carbides at a depth of 1/10 of the sheet thickness from the surfaces of the steel sheets. Fig. 1 shows a sample to which Sb was added and which was cooled under Condition (e), Fig. 2 shows a sample (Table 3, column c, bottom) to which Sb was added and which was cooled under Condition (c), Fig. 3 shows a sample which had no additive component and which was cooled under Condition (e) (Table 3, column 3, top), and Fig. 4 shows a sample which had no additive component and which was cooled under Condition (c) (Table 3, column c, top).
As is shown in Table 3, the magnetic flux density (B8)(T) of the sample to which Sb was added (bottom half of Table 3) and which was manufactured under the intermediate annealing cooling condition (c) (Table 3, column c) was particularly high. In this sample, carbide precipitates having a size ranging from 300 to 500 A and sparsely precipitated were observed after the intermediate annealing, and are shown in Fig. 2, as heretofore noted. In contrast, in the sample which had no additive component and which was manufactured under the same cooling condition (c) (Table 3, column c, top), fine carbide precipitates having a size of about 100 A were undesirably precipitated at a high density, as shown in ~ig.
4.
With respect to the steel sheets which had no additive - component, in the case of creating a work strain by skin-pass-rolling in accordance with the condition (c), carbide precipitation sites were increased during cooling so that carbides were finely precipitated at a high density, as is apparent from comparison with processing under the condition (b). In contrast, with respect to the steel sheets to which Sb was added, precipitation sites were not increased and slightly coarse precipitates were observed. According to our study after these experiments, such sparse precipitation of carbides having a size ranging from 300 to 500 ~ increases the (111) intensity of the structure primarily recrystallized by decarburization annealing after final cold rolling and reduces the {lll} <uvw>, in particular the {111} <110> intensity while increasing the {111} <112> intensity. The {111} <110> grains limit the growth of the (110) [001] secondary grains which contribute to the increase in the magnetic flux density, while the {111} <112> grains promote the growth of (110) [001]
secondary grains. It is thought that addition of Sb in the particular process provides this effect and enables formation of a product having a substantially high magnetic flux density as in the case of Condition (c) as shown in the top portion of Table 3.
It is thought that this effect of Sb in steel relates to segregation of Sb, that Sb is segregated at base points in crystal grains such as to form carbide precipitation sites, and that this segregation results from the limitation of precipitation carbides during cooling.
This action of Sb is particularly effective in a temperature range of about 200 to 500C; the amount of strain to be applied may be very small, e.g., about 0.005 to 3 %.
2~0i~
It has also been found that the aging effect at the time of final cold rolling can also be improved according to this invention because the amount of solid solution carbon is increased by the carbide precipitation limiting effect of Sb.
It is known that a small strain of 0.5 % created by skin-pass-rolling is concentrated at a surface-layer portion of the steel sheet. In this work as well, the form of precipitated carbides was changed according to the change in the amount of strain in the thickness direction of the sheet, and the density of precipitated carbides was reduced toward the center of the sheet in the thickness direction.
The fact that the form of precipitated carbides was changed in the sheet thickness direction is regarded as a reason for the success of this work. To positively utilize this effect, a similar experiment was also conducted by creating a strain of 0.5 % by bending with a leveler, and suitable effects were thereby obtained.
A carbide precipitation processing method is disclosed in Japanese Patent Laid-Open 61-149432. In this method, high-density dislocations uniform in the direction of sheet thickness are provided by rolling at a high temperature of 1,000 to 400C, and the speed of cooling in a step of precipitating carbon is high, such as 10C/s. This method is intended to precipitate finely divided carbides and to increase the (110) [001] intensity of the texture of the product.
Japanese Patent Laid-Open 58-15797 also discloses a technique for precipitating carbides of a size of 100 to 500 ~. In this case, however, the precipitation temperature range is a range of low temperatures, i.e., 300 to 150C, and the effect of Sb is not effectively utilized, and there is no disclosure or suggestion of our special ideas relating to the precipitation processing which constitutes a feature of the present invention, including that of creating a strain during precipitation. This technique is therefore sharply different from the present invention with respect to the carbide precipitation density and requires high-density precipitation for increasing the (110) [001] intensity as in the case of the method disclosed in Japanese Patent Laid-Open 61-149432.
In contrast, in accordance with the present invention, it is important to precipitate carbides sparsely to reduce the {111} <uvw> intensity, in particular the {111} <110> intensity of the primary recrystallized structure while increasing its {111} <112> intensity.
It is important to define the ranges of chemical components of the composition of the oriented silicon steel sheet in accordance with the present invention. Preferable ranges of the components will be described below.
C is necessary for improving the hot-rolled structure of the steel. However, if the C content is excessive, it is difficult to decarburize the steel. It is therefore preferable to limit the carbon content to a range of about 0.035 to 0.090 % by weight.
If the Si content is below a lower limit the desired core loss characteristic cannot be obtained. If the Si content is excessive it is difficult to perform cold rolling. It is preferable to provide an Si content in the range of about 2.5 to 4.5 % by weight.
Mn can be utilized as an inhibitor component. In case of an excessively large amount of Mn, Mn compound in the steel cannot be dissolved during slab-reheating process, and it is accordingly preferable to provide an Mn content in the range of about 0.05 to 0.15 % by weight.
S or Se is effective when combined with Mn to form MnS
or MnSe which acts as an inhibitor. The range of S or Se content for finely precipitating MnS or MnSe is preferably about 0.01 to 0.04 % by weight in either case of whether used alone or together.
It is specifically necessary for the steel sheet of the present invention to contain acid-soluble Al or N as inhibitor components for the purpose of achieving a high magnetic flux density, and addition of certain amounts of acid-soluble Al or N is required. However, if these contents are excessive fine precipitation is difficult. It is preferable to maintain the content of acid-soluble Al to a range of about 0.01 to 0.15 % by weight and the content of N to a range of about 0.0030 to 0.020 % by weight.
Further, according to the present invention, the presence of Sb in the steel is indispensable, and it is possible to limit precipitation of C at grain boundaries or in crystal grains in the steel by providing a content of Sb. To enable such an effect, about 0.005 % or greater by weight of Sb is necessary. However, if the Sb content exceeds about 0.040 %
by weight, the problem of grain boundary embrittlement is encountered, and it is difficult to perform cold rolling. The Sb content is therefore maintained within a range of about 0.005 to 0.040 % by weight.
To improve magnetic properties, other inhibitor strengthening components such as Cu, Cr, Bi, Sn, B, Ge and the like may be added as desired. The content of each of such components may be within well-known ranges. To prevent occurrence of surface defects due to hot-rolling embrittlement, it is preferable to add Mo in a range of about 0.005 to 0.020 % by weight.
Next, a process of manufacture in accordance with the present invention will be described below.
Well-known manufacturing methods are applied for manufacturing the steel sheet, and ingots or slabs are reproduced as desired, adjusted to the desired size, and thereafter heated and hot-rolled. The hot-rolled steel sheet is processed by cold rolling one time or in a plurality of stages until its thickness is reduced to a desired final thickness.
For annealing before final cold rolling a high temperature in a range of about 850 to 1,200C is required to dissolve AlN, and, after this annealing, quenching to 500C
or lower is required to precipitate AlN and it is also necessary to prevent precipitation of C at grain boundaries.
If the cooling speed is lower than 15C/s, C is pxecipitated at grain boundaries, or, if the cooling speed exceeds 500C/s, the shape of the steel sheet after the cooling is deteriorated. The cooling rate is therefore maintained within a range of about 15 to 500C/s.
Thereafter, a small strain ranglng from about 0.005 to 3.0 % is created ln a temperature range from the temperature reached by ~uenching (about 500C at the maximum) to about 200C. The steel sheet is cooled slowly at a coolln~
speed of about 2C/s or lower during this straining or after the straining for a perlod of tlme of about 60 to 1~0 seconds ln which the steel sheet is maintained in essentially the same temperature range.
This step ls intended to preclpitate sparsely arranged carbldes havlng a slze ranging from about 300 to 500 ~ ln grains, whlch effect relates to one of the most lmportant features of the present lnvention. This processinq ls performed withln a hlgh temperature range from the temperature reached by cooling, i.e., about 500C at the maximum to about 200C, and a strain is created in this temperature range, a feature unknown before the present invention. The precipitation of carbides is controlled to provide the deslred slze and denslty by balanclng three influenclng factors including (a) the fact that the C
diffusion speed is comparatively high so that carbides are coarsely formed, (b) the fact that the carbide precipitation polnts are lncreased by straining so that carbides precipltate flnely at a hlgh density, and (c) the fact that precipitation of carbides at graln boundarles and ln crystal gralns ls llmlted by the segregation effect of the presence of Sb.
E
73461-~7 , . . _ .. ., .. ~
Carbide preclpltates have an excesslvely large slze if the precipitation temperature exceeds about 500C. They are ~, 23a 204801~
excessively fine if the precipitation temperature is lower than about 200C. Preferably the temperature at which precipitation is performed is within the range of about 450C
to 300C.
If the maintenance time is shorter than about 60 seconds, the carbides are not formed sufficiently coarsely. If it is longer than about 180 seconds, carbides are formed excessively coarsely, and the number of precipitation points is increased and the amount of solid solution is considerably reduced, with undesirable results.
When slow cooling is performed instead of the constant-temperature maintenance step it is necessary to set the cooling speed to about 2C/s or lower.
It is necessary to effect straining immediately after quenching or in the temperature range of about 500 to 200C
before the carbon precipitation processing. It is thereby possible to prevent carbides from precipitating excessively coarsely. If the amount of strain provided is less than about 0.005 % by weight, the carbides are formed excessively coarsely. If the strain is more than about 3.0 %, carbides are finely precipitated at an excessively high density. The amount of strain is therefore set within a range of about 0.005 to 3.0 ~. A range of 0.01 to 1.0 % is particularly preferable.
Needless to say, straining may be performed by any conventional straining method, e.g., a skin pass method based on rolling, a bending method using a bending roll, a straining method using a leveler roll, shot blasting, or the like.
20~8014 The steel sheet is then subjected to final cold rolling.
At this time, to obtain a high magnetic flux density, it is necessary to set the rolling reduction to a range of about 80 to 95 %, as is well known.
Performing well-known aging or hot rolling treatment during this final cold rolling is further effective in the process of the present invention, because the amount of solid solution C in the steel of the present invention is large.
The aging temperature is preferably adjusted to the range of about 200 to 400C. If the aging temperature is higher than about 400C the shapes of precipitated carbides are changed so that the object of the present invention cannot be achieved. If the aging temperature is lower than about 200C, solid solution C or solid solution N is not sufficiently fixed on dislocations, and further improvements in characteristics cannot be expected.
It is necessary to set the rolling reduction to a range of about 80 to 95 %, as is well known. If the rolling reduction is less than about 80 %, a sufficiently high magnetic flux density cannot be obtained. If the rolling reduction exceeds about 95 %, it is difficult to develop secondary recrystallization grains.
The steel sheet after final cold rolling is degreased and is then annealed for decarburization and primary recrystallization. An annealing separation agent having MgO
as a main component is thereafter applied to the steel sheet, and the steel sheet is coiled to be subjected to finishing annealing and is coated with an insulating material if 204801~
necessary. Needless to say, the steel sheet may also be processed to fractionate magnetic domains by laser, plasma or any other means.
(Examples) Example 1 Eleven steel ingots B, D, E, F, G, H, I, J, X, L, and M
shown in Table 4 were provided in conformity with the present invention. These steels and other two steels A, C provided as comparative examples, thirteen steels in all were hot rolled in a conventional manner to form hot-rolled coils each having a thickness of 2.2 mm.
20'18014 n ~ I ~ ~
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n 204801~
Each steel sheet was thereafter subjected to normal annealing at l,000C for 90 seconds and was cold-rolled until its thickness was reduced to an intermediate thickness of 1.50 mm. The reduced steel sheet was further annealed at 1,100C
for 90 seconds, quenched at a rate of 60C/s to 350C, and passed through a slow cooling box having a bending roll and was thereby strained to an extent of 1.5 ~ while being cooled at a rate of 2C/s to 200C. The steel sheet was thereafter cooled in atmospheric air.
- The steel sheet was then rolled until its thickness was reduced to a final thickness of 0.22 mm, electrolytically degreased, and subjected to decarburization/primary recrystallization annealing at 850C for 2 minutes in a wet hydrogen atmosphere. An MgO agent containing 5 % TiO2 was then applied to the steel sheet, and the steel sheet was subjected to finishing annealing at 1,200C for 10 hours.
Thereafter, the surfaces of the sheet were coated to give the steel sheet tensile stress and were partially processed to fractionate magnetic domains at 10 mm pitches by the plasma jet method. Table 5 shows the magnetic characteristics before and after the magnetic domain fractionating processing of the steel sheets.
204801~
Table 5 Magnetic domain MagnetiC flux Core loss Ingot fractionating density Wl7/so (W/kg)Note symbol processing B8 (T) A Unprocessed 1.875 1.15 Comparative Processed 1.874 1.09 example B -Unprocessed 1.935 0.92 Conformable Processed 1.936 0.78 example C Unprocessed 1.883 1.07 Comparative Processed 1.883 1.02 example D Unprocessed 1.938 0.95 Conformable Processed 1.938 0.84 example E Unprocessed 1.941 0.87 ditto Processed 1.942 0.73 F Unprocessed 1.946 0.85 ditto Processed 1.945 0.70 G Unprocessed 1.942 0.86 ditto Processed 1.943 0.72 H Unprocessed 1.937 0.97 ditto Processed 1.938 0.83 I Unprocessed 1.940 0.87 ditto Processed 1.941 0.72 J Unprocessed 1.941 0.83 ditto Processed 1.941 0.70 K Unprocessed 1.938 0.86 ditto Processed 1.937 0.73 L Unprocessed 1.942 0.85 ditto Processèd 1.943 0.71 M Unprocessed 1.939 0.88 ditto Processed 1.938 0.75 Note: * Magnetic domain fractionating at lOmm pitches by plasma jet method 20~8~4 As appears in Table 5, the conformable examples (all except A and C) have characteristics improved in magnetic flux density and core loss due to this invention, in comparison with those of the comparative Examples A and C. The magnetic flux density of the conformable examples was 1.946 T (Ingot F) at the m~;mum with respect to B8, as compared to 1.875 and 1.883 for comparative Examples A and C. The magnetic domain fractionating processing remarkably improved the core loss but did not substantially adversely influence the magnetic flux density.
Example 2 The steel ingot F shown in Table 4 was hot-rolled in a conventional manner to provide hot-rolled steel sheets having thicknesses of 2.4, 2.2, 2.0, and 1.5 mm.
The hot-rolled steel sheets having thicknesses of 2.4 and 2.2 mm were respectively annealed at 1,175C for 90 seconds and at 1,150C for 90 seconds, then quenched to 400C at an average cooling speed of 50C/s, strained to an extent of 2 % by a hot skin pass roller, slowly cooled to 250C at an average cooling speed of 1.5C/s, and quenched in water.
Thereafter, these steel sheets were respectively cold-rolled to final thicknesses of 0.30 and 0.28 mm. When the thicknesses of these steel sheets were respectively reduced to 1.3 and 1.0 mm, each sheet was separated into two. One of them was successively cold-rolled and the other was aged at 300C for 2 minutes and cold-rolled to the final thickness.
The hot-rolled steel sheets having thicknesses of 2.0 and 1.5 mm were normalized at 1,000C for 90 seconds, naturally 20~801~
cooled, respectively cold-rolled to thicknesses of 1.4 and 1.1 mm, annealed at 1,100C for 90 seconds, and quenched to 350C
at an average speed of 60C/s. They were then strained to an extent of 1.0 % by a hot leveler, maintained at 320C for 120 seconds, and taken out of the furnace and naturally cooled.
Thereafter they were respectively cold-rolled to final thicknesses of 0.20 and 0.15 mm. When the thicknesses of these steel sheets were respectively reduced to 0.7 and 0.55 mm, each sheet was separated into two. One of them was successively cold-rolled and the other was aged at 300C for 2 minutes and cold-rolled to the final thickness. After final cold rolling the steel sheets were degreased and subjected to decarburization/primary recrystallization annealing at 850C
for 2 minutes in a wet hydrogen atmosphere. An MgO containing 2 % SrSO4 was then applied to the steel sheets, and the steel sheets were subjected to finishing annealing at 1,200C for 10 hours. Thereafter the surfaces of the sheets were coated to give a tensile stress to the sheets and processed to fractionate magnetic domains by 5 mm pitch electron beam irradiation. Table 6 shows the magnetic characteristics of the steel sheets thus processed.
Table 6 Item Non-aged Aged *
Final thickness (mm) B8 (T)Wl7/so (W~Kg) B8 (T)W17/50 (W/Kg) 0.30 1.9420.97 1.9450.90 0.28 1.9480.93 1.9440.88 0.20 1.9400.87 1.9420.82 w N
0.15 1.934 0.86 1.930 0.77 Note: * Aged at 300C for 2 minutes during cold rolling o 204801~
As appears in Table 6 the magnetic flux density was improved even though the final thickness was substantially reduced down to 0.15 mm, and the magnetic domain fractionating processing during the cold rolling remarkably improved the core loss but did not substantially influence the magnetic flux density.
Example 3 The ingot G shown in Table 4 was hot-rolled in a conventional manner to provide a hot-rolled coil having a thickness of 2.0 mm. This steel sheet was normalized at - 1,000C for 90 seconds and was cold-rolled to an intermediate thickness of 1.50 mm. ThiS steel sheet was separated into three pieces and all were subjected to intermediate annealing at l,100C for 90 seconds. This cooling was performed under three different sets of conditions.
The first set of conditions (I) was that the steel sheet was cooled in hot water at 80C.
The second set of conditions (II) was that the steel sheet was cooled to 350C at an average cooling speed of 60C/s, was slowly cooled to 300C for 2 minutes while being strained to an extent of 0.5 % by a bending roll, and was cooled in atmospheric air.
The third set of conditions (III) was that the steel sheet was cooled to 400C at an average cooling speed of 60C/s, was cooled to 250C at a cooling speed of 2C/s, and was cooled in atmospheric air.
Each of these three steel sheets was separated into two.
One of them was cold-rolled in a conventional manner to a 2~801~
final thickness of 0.20 mm, while the other was hot-rolled at 250C to a final thickness of 0.20 mm. After final cold rolling, all the steel sheets were degreased and subjected to decarburization/primary recrystallization annealing at 860C
for 2 minutes in a wet hydrogen atmosphere. An MgO separator containing 10 % TiOz was then applied to the steel sheets, and the steel sheets were subjected to finishing annealing at 1,200C for 10 hours. Thereafter the surfaces of the sheets were tension-coated and the magnetic characteristics were measured. Table 7 shows the results of this measurement.
Table 7 Item Normally rolled sheet Warm-rolled sheet *
Cooling Note conditionB8 (T) Wl7/50 (W/Kg) B8 (T) Wl7/so (W/Kg) (I) 1.882 1.08 1.888 0.97 Comparative example (II) 1.939 0.85 1.941 0.82 Conformable example w (III) 1.896 1.05 1.894 1.95 Comparative example Note: * Finishing-cold-rolled at 250C
o 2~01~
As shown in Table 7, the conformable example processed under the cooling conditions (II) was improvçd in both magnetic flux density and core loss in comparison with the comparative examples processed under the cooling conditions (I) and (III), and it was found that the creation of a small strain in a temperature range of 500 to 200C during the cooling for the annealing before the final cold rolling was effective in improving the magnetic characteristics of the sheet.
According to the present invention, a silicon steel sheet cont~ining Al and Sb is used and cooling control and creation of a small strain are effected during cooling for annealing before final cold rolling, so that an oriented silicon steel sheet having a high magnetic flux density can be manufactured with stability even if the sheet thickness is reduced. The oriented silicon steel sheet manufactured in accordance with the present invention has excellent properties for use in transformer cores and other products having high magnetic flux density and good stability with reduced core loss.
It has also been found that the aging effect at the time of final cold rolling can also be improved according to this invention because the amount of solid solution carbon is increased by the carbide precipitation limiting effect of Sb.
It is known that a small strain of 0.5 % created by skin-pass-rolling is concentrated at a surface-layer portion of the steel sheet. In this work as well, the form of precipitated carbides was changed according to the change in the amount of strain in the thickness direction of the sheet, and the density of precipitated carbides was reduced toward the center of the sheet in the thickness direction.
The fact that the form of precipitated carbides was changed in the sheet thickness direction is regarded as a reason for the success of this work. To positively utilize this effect, a similar experiment was also conducted by creating a strain of 0.5 % by bending with a leveler, and suitable effects were thereby obtained.
A carbide precipitation processing method is disclosed in Japanese Patent Laid-Open 61-149432. In this method, high-density dislocations uniform in the direction of sheet thickness are provided by rolling at a high temperature of 1,000 to 400C, and the speed of cooling in a step of precipitating carbon is high, such as 10C/s. This method is intended to precipitate finely divided carbides and to increase the (110) [001] intensity of the texture of the product.
Japanese Patent Laid-Open 58-15797 also discloses a technique for precipitating carbides of a size of 100 to 500 ~. In this case, however, the precipitation temperature range is a range of low temperatures, i.e., 300 to 150C, and the effect of Sb is not effectively utilized, and there is no disclosure or suggestion of our special ideas relating to the precipitation processing which constitutes a feature of the present invention, including that of creating a strain during precipitation. This technique is therefore sharply different from the present invention with respect to the carbide precipitation density and requires high-density precipitation for increasing the (110) [001] intensity as in the case of the method disclosed in Japanese Patent Laid-Open 61-149432.
In contrast, in accordance with the present invention, it is important to precipitate carbides sparsely to reduce the {111} <uvw> intensity, in particular the {111} <110> intensity of the primary recrystallized structure while increasing its {111} <112> intensity.
It is important to define the ranges of chemical components of the composition of the oriented silicon steel sheet in accordance with the present invention. Preferable ranges of the components will be described below.
C is necessary for improving the hot-rolled structure of the steel. However, if the C content is excessive, it is difficult to decarburize the steel. It is therefore preferable to limit the carbon content to a range of about 0.035 to 0.090 % by weight.
If the Si content is below a lower limit the desired core loss characteristic cannot be obtained. If the Si content is excessive it is difficult to perform cold rolling. It is preferable to provide an Si content in the range of about 2.5 to 4.5 % by weight.
Mn can be utilized as an inhibitor component. In case of an excessively large amount of Mn, Mn compound in the steel cannot be dissolved during slab-reheating process, and it is accordingly preferable to provide an Mn content in the range of about 0.05 to 0.15 % by weight.
S or Se is effective when combined with Mn to form MnS
or MnSe which acts as an inhibitor. The range of S or Se content for finely precipitating MnS or MnSe is preferably about 0.01 to 0.04 % by weight in either case of whether used alone or together.
It is specifically necessary for the steel sheet of the present invention to contain acid-soluble Al or N as inhibitor components for the purpose of achieving a high magnetic flux density, and addition of certain amounts of acid-soluble Al or N is required. However, if these contents are excessive fine precipitation is difficult. It is preferable to maintain the content of acid-soluble Al to a range of about 0.01 to 0.15 % by weight and the content of N to a range of about 0.0030 to 0.020 % by weight.
Further, according to the present invention, the presence of Sb in the steel is indispensable, and it is possible to limit precipitation of C at grain boundaries or in crystal grains in the steel by providing a content of Sb. To enable such an effect, about 0.005 % or greater by weight of Sb is necessary. However, if the Sb content exceeds about 0.040 %
by weight, the problem of grain boundary embrittlement is encountered, and it is difficult to perform cold rolling. The Sb content is therefore maintained within a range of about 0.005 to 0.040 % by weight.
To improve magnetic properties, other inhibitor strengthening components such as Cu, Cr, Bi, Sn, B, Ge and the like may be added as desired. The content of each of such components may be within well-known ranges. To prevent occurrence of surface defects due to hot-rolling embrittlement, it is preferable to add Mo in a range of about 0.005 to 0.020 % by weight.
Next, a process of manufacture in accordance with the present invention will be described below.
Well-known manufacturing methods are applied for manufacturing the steel sheet, and ingots or slabs are reproduced as desired, adjusted to the desired size, and thereafter heated and hot-rolled. The hot-rolled steel sheet is processed by cold rolling one time or in a plurality of stages until its thickness is reduced to a desired final thickness.
For annealing before final cold rolling a high temperature in a range of about 850 to 1,200C is required to dissolve AlN, and, after this annealing, quenching to 500C
or lower is required to precipitate AlN and it is also necessary to prevent precipitation of C at grain boundaries.
If the cooling speed is lower than 15C/s, C is pxecipitated at grain boundaries, or, if the cooling speed exceeds 500C/s, the shape of the steel sheet after the cooling is deteriorated. The cooling rate is therefore maintained within a range of about 15 to 500C/s.
Thereafter, a small strain ranglng from about 0.005 to 3.0 % is created ln a temperature range from the temperature reached by ~uenching (about 500C at the maximum) to about 200C. The steel sheet is cooled slowly at a coolln~
speed of about 2C/s or lower during this straining or after the straining for a perlod of tlme of about 60 to 1~0 seconds ln which the steel sheet is maintained in essentially the same temperature range.
This step ls intended to preclpitate sparsely arranged carbldes havlng a slze ranging from about 300 to 500 ~ ln grains, whlch effect relates to one of the most lmportant features of the present lnvention. This processinq ls performed withln a hlgh temperature range from the temperature reached by cooling, i.e., about 500C at the maximum to about 200C, and a strain is created in this temperature range, a feature unknown before the present invention. The precipitation of carbides is controlled to provide the deslred slze and denslty by balanclng three influenclng factors including (a) the fact that the C
diffusion speed is comparatively high so that carbides are coarsely formed, (b) the fact that the carbide precipitation polnts are lncreased by straining so that carbides precipltate flnely at a hlgh density, and (c) the fact that precipitation of carbides at graln boundarles and ln crystal gralns ls llmlted by the segregation effect of the presence of Sb.
E
73461-~7 , . . _ .. ., .. ~
Carbide preclpltates have an excesslvely large slze if the precipitation temperature exceeds about 500C. They are ~, 23a 204801~
excessively fine if the precipitation temperature is lower than about 200C. Preferably the temperature at which precipitation is performed is within the range of about 450C
to 300C.
If the maintenance time is shorter than about 60 seconds, the carbides are not formed sufficiently coarsely. If it is longer than about 180 seconds, carbides are formed excessively coarsely, and the number of precipitation points is increased and the amount of solid solution is considerably reduced, with undesirable results.
When slow cooling is performed instead of the constant-temperature maintenance step it is necessary to set the cooling speed to about 2C/s or lower.
It is necessary to effect straining immediately after quenching or in the temperature range of about 500 to 200C
before the carbon precipitation processing. It is thereby possible to prevent carbides from precipitating excessively coarsely. If the amount of strain provided is less than about 0.005 % by weight, the carbides are formed excessively coarsely. If the strain is more than about 3.0 %, carbides are finely precipitated at an excessively high density. The amount of strain is therefore set within a range of about 0.005 to 3.0 ~. A range of 0.01 to 1.0 % is particularly preferable.
Needless to say, straining may be performed by any conventional straining method, e.g., a skin pass method based on rolling, a bending method using a bending roll, a straining method using a leveler roll, shot blasting, or the like.
20~8014 The steel sheet is then subjected to final cold rolling.
At this time, to obtain a high magnetic flux density, it is necessary to set the rolling reduction to a range of about 80 to 95 %, as is well known.
Performing well-known aging or hot rolling treatment during this final cold rolling is further effective in the process of the present invention, because the amount of solid solution C in the steel of the present invention is large.
The aging temperature is preferably adjusted to the range of about 200 to 400C. If the aging temperature is higher than about 400C the shapes of precipitated carbides are changed so that the object of the present invention cannot be achieved. If the aging temperature is lower than about 200C, solid solution C or solid solution N is not sufficiently fixed on dislocations, and further improvements in characteristics cannot be expected.
It is necessary to set the rolling reduction to a range of about 80 to 95 %, as is well known. If the rolling reduction is less than about 80 %, a sufficiently high magnetic flux density cannot be obtained. If the rolling reduction exceeds about 95 %, it is difficult to develop secondary recrystallization grains.
The steel sheet after final cold rolling is degreased and is then annealed for decarburization and primary recrystallization. An annealing separation agent having MgO
as a main component is thereafter applied to the steel sheet, and the steel sheet is coiled to be subjected to finishing annealing and is coated with an insulating material if 204801~
necessary. Needless to say, the steel sheet may also be processed to fractionate magnetic domains by laser, plasma or any other means.
(Examples) Example 1 Eleven steel ingots B, D, E, F, G, H, I, J, X, L, and M
shown in Table 4 were provided in conformity with the present invention. These steels and other two steels A, C provided as comparative examples, thirteen steels in all were hot rolled in a conventional manner to form hot-rolled coils each having a thickness of 2.2 mm.
20'18014 n ~ I ~ ~
Z v ~ D ' V ~ D ,_, ,,, ,,, ,V ,V V ~ V v z ~ ~ U~ ~ O ~ ~ ~ U~ ~ o ~
-6 ~ ~ ~ ~ ~ N ~ ~ ~
-O
V V V V V V V VV V J- .
o N N ~N
eo o o o o oo oo o ~o o U~ . .. .. . .. . .
o o o o o oo oo o oo o ~O O O O O OO OO O OO O
U
O O O O O OO OO O OO O
In V V V V V VV VV . VV V
o O O _I~ `1 ~ ~'1 t~
rn V
o o o oo o o o o o o o :~ O O O OO O O O O O O O O
o O O O O OO O O O O O O O
~r E- v. o ~ ~ ~ ~~ ~ o ~V V V V VV . VVVVV~
O
r a~ oa~co o N O ~ ~ a~
rn v v v o o oo o o o o o o _Io o o oo o o o o o o o rn o o o oo o o o o o o o o .
o o o o oo o o o o o o o ,~ O O O OO O O O O O O O O
O O O O OO O O O O O O O
O O O OO O O
O O O O OO O O O O O O O
O O O O OO O O O O O O O
.
O O O O OO O O O O O O O
rn o~ u~ C O ~ ~ ~ ~
u o o o o oo o o o o o o o - o o o o oo o o o o o o o o ~
r~ u a ~~ r~
n 204801~
Each steel sheet was thereafter subjected to normal annealing at l,000C for 90 seconds and was cold-rolled until its thickness was reduced to an intermediate thickness of 1.50 mm. The reduced steel sheet was further annealed at 1,100C
for 90 seconds, quenched at a rate of 60C/s to 350C, and passed through a slow cooling box having a bending roll and was thereby strained to an extent of 1.5 ~ while being cooled at a rate of 2C/s to 200C. The steel sheet was thereafter cooled in atmospheric air.
- The steel sheet was then rolled until its thickness was reduced to a final thickness of 0.22 mm, electrolytically degreased, and subjected to decarburization/primary recrystallization annealing at 850C for 2 minutes in a wet hydrogen atmosphere. An MgO agent containing 5 % TiO2 was then applied to the steel sheet, and the steel sheet was subjected to finishing annealing at 1,200C for 10 hours.
Thereafter, the surfaces of the sheet were coated to give the steel sheet tensile stress and were partially processed to fractionate magnetic domains at 10 mm pitches by the plasma jet method. Table 5 shows the magnetic characteristics before and after the magnetic domain fractionating processing of the steel sheets.
204801~
Table 5 Magnetic domain MagnetiC flux Core loss Ingot fractionating density Wl7/so (W/kg)Note symbol processing B8 (T) A Unprocessed 1.875 1.15 Comparative Processed 1.874 1.09 example B -Unprocessed 1.935 0.92 Conformable Processed 1.936 0.78 example C Unprocessed 1.883 1.07 Comparative Processed 1.883 1.02 example D Unprocessed 1.938 0.95 Conformable Processed 1.938 0.84 example E Unprocessed 1.941 0.87 ditto Processed 1.942 0.73 F Unprocessed 1.946 0.85 ditto Processed 1.945 0.70 G Unprocessed 1.942 0.86 ditto Processed 1.943 0.72 H Unprocessed 1.937 0.97 ditto Processed 1.938 0.83 I Unprocessed 1.940 0.87 ditto Processed 1.941 0.72 J Unprocessed 1.941 0.83 ditto Processed 1.941 0.70 K Unprocessed 1.938 0.86 ditto Processed 1.937 0.73 L Unprocessed 1.942 0.85 ditto Processèd 1.943 0.71 M Unprocessed 1.939 0.88 ditto Processed 1.938 0.75 Note: * Magnetic domain fractionating at lOmm pitches by plasma jet method 20~8~4 As appears in Table 5, the conformable examples (all except A and C) have characteristics improved in magnetic flux density and core loss due to this invention, in comparison with those of the comparative Examples A and C. The magnetic flux density of the conformable examples was 1.946 T (Ingot F) at the m~;mum with respect to B8, as compared to 1.875 and 1.883 for comparative Examples A and C. The magnetic domain fractionating processing remarkably improved the core loss but did not substantially adversely influence the magnetic flux density.
Example 2 The steel ingot F shown in Table 4 was hot-rolled in a conventional manner to provide hot-rolled steel sheets having thicknesses of 2.4, 2.2, 2.0, and 1.5 mm.
The hot-rolled steel sheets having thicknesses of 2.4 and 2.2 mm were respectively annealed at 1,175C for 90 seconds and at 1,150C for 90 seconds, then quenched to 400C at an average cooling speed of 50C/s, strained to an extent of 2 % by a hot skin pass roller, slowly cooled to 250C at an average cooling speed of 1.5C/s, and quenched in water.
Thereafter, these steel sheets were respectively cold-rolled to final thicknesses of 0.30 and 0.28 mm. When the thicknesses of these steel sheets were respectively reduced to 1.3 and 1.0 mm, each sheet was separated into two. One of them was successively cold-rolled and the other was aged at 300C for 2 minutes and cold-rolled to the final thickness.
The hot-rolled steel sheets having thicknesses of 2.0 and 1.5 mm were normalized at 1,000C for 90 seconds, naturally 20~801~
cooled, respectively cold-rolled to thicknesses of 1.4 and 1.1 mm, annealed at 1,100C for 90 seconds, and quenched to 350C
at an average speed of 60C/s. They were then strained to an extent of 1.0 % by a hot leveler, maintained at 320C for 120 seconds, and taken out of the furnace and naturally cooled.
Thereafter they were respectively cold-rolled to final thicknesses of 0.20 and 0.15 mm. When the thicknesses of these steel sheets were respectively reduced to 0.7 and 0.55 mm, each sheet was separated into two. One of them was successively cold-rolled and the other was aged at 300C for 2 minutes and cold-rolled to the final thickness. After final cold rolling the steel sheets were degreased and subjected to decarburization/primary recrystallization annealing at 850C
for 2 minutes in a wet hydrogen atmosphere. An MgO containing 2 % SrSO4 was then applied to the steel sheets, and the steel sheets were subjected to finishing annealing at 1,200C for 10 hours. Thereafter the surfaces of the sheets were coated to give a tensile stress to the sheets and processed to fractionate magnetic domains by 5 mm pitch electron beam irradiation. Table 6 shows the magnetic characteristics of the steel sheets thus processed.
Table 6 Item Non-aged Aged *
Final thickness (mm) B8 (T)Wl7/so (W~Kg) B8 (T)W17/50 (W/Kg) 0.30 1.9420.97 1.9450.90 0.28 1.9480.93 1.9440.88 0.20 1.9400.87 1.9420.82 w N
0.15 1.934 0.86 1.930 0.77 Note: * Aged at 300C for 2 minutes during cold rolling o 204801~
As appears in Table 6 the magnetic flux density was improved even though the final thickness was substantially reduced down to 0.15 mm, and the magnetic domain fractionating processing during the cold rolling remarkably improved the core loss but did not substantially influence the magnetic flux density.
Example 3 The ingot G shown in Table 4 was hot-rolled in a conventional manner to provide a hot-rolled coil having a thickness of 2.0 mm. This steel sheet was normalized at - 1,000C for 90 seconds and was cold-rolled to an intermediate thickness of 1.50 mm. ThiS steel sheet was separated into three pieces and all were subjected to intermediate annealing at l,100C for 90 seconds. This cooling was performed under three different sets of conditions.
The first set of conditions (I) was that the steel sheet was cooled in hot water at 80C.
The second set of conditions (II) was that the steel sheet was cooled to 350C at an average cooling speed of 60C/s, was slowly cooled to 300C for 2 minutes while being strained to an extent of 0.5 % by a bending roll, and was cooled in atmospheric air.
The third set of conditions (III) was that the steel sheet was cooled to 400C at an average cooling speed of 60C/s, was cooled to 250C at a cooling speed of 2C/s, and was cooled in atmospheric air.
Each of these three steel sheets was separated into two.
One of them was cold-rolled in a conventional manner to a 2~801~
final thickness of 0.20 mm, while the other was hot-rolled at 250C to a final thickness of 0.20 mm. After final cold rolling, all the steel sheets were degreased and subjected to decarburization/primary recrystallization annealing at 860C
for 2 minutes in a wet hydrogen atmosphere. An MgO separator containing 10 % TiOz was then applied to the steel sheets, and the steel sheets were subjected to finishing annealing at 1,200C for 10 hours. Thereafter the surfaces of the sheets were tension-coated and the magnetic characteristics were measured. Table 7 shows the results of this measurement.
Table 7 Item Normally rolled sheet Warm-rolled sheet *
Cooling Note conditionB8 (T) Wl7/50 (W/Kg) B8 (T) Wl7/so (W/Kg) (I) 1.882 1.08 1.888 0.97 Comparative example (II) 1.939 0.85 1.941 0.82 Conformable example w (III) 1.896 1.05 1.894 1.95 Comparative example Note: * Finishing-cold-rolled at 250C
o 2~01~
As shown in Table 7, the conformable example processed under the cooling conditions (II) was improvçd in both magnetic flux density and core loss in comparison with the comparative examples processed under the cooling conditions (I) and (III), and it was found that the creation of a small strain in a temperature range of 500 to 200C during the cooling for the annealing before the final cold rolling was effective in improving the magnetic characteristics of the sheet.
According to the present invention, a silicon steel sheet cont~ining Al and Sb is used and cooling control and creation of a small strain are effected during cooling for annealing before final cold rolling, so that an oriented silicon steel sheet having a high magnetic flux density can be manufactured with stability even if the sheet thickness is reduced. The oriented silicon steel sheet manufactured in accordance with the present invention has excellent properties for use in transformer cores and other products having high magnetic flux density and good stability with reduced core loss.
Claims (10)
1. A method of manufacturing an oriented silicon steel sheet having improved magnetic characteristics, which method comprises:
(a) providing a hot-rolled steel sheet of a silicon steel having a composition containing about 0.01 to 0.15 % by weight of acid-soluble Al, about 0.005 to 0.04 % by weight of Sb, about 0.035 to 0.090 % by weight of C, about 2.5 to 4.5 %
by weight of Si, about 0.05 to 0.15 % by weight of Mn, 0.01 to 0.04 % by weight of S or Se and the balance being essentially Fe and inevitable impurities;
(b) softening-annealing the hot-rolled steel sheet at a temperature of about 850 to 1,200°C to dissolve AlN
before final cold rolling;
(c) successively quenching the softening-annealed steel sheet at a cooling speed of about 15 to 500°C/s to a temperature of about 500°C or less, thereby precipitating AlN
without precipitating C at grain boundaries;
(d) applying to the quenched steel sheet a strain ranging from about 0.005 to 3.0 % while maintaining the sheet at a temperature in the range from about the temperature reached by quenching to about 200°C;
(e) controlling carbide precipitation at an effective cooling speed of about 2°C/s or less by cooling the steel sheet during the straining step (d) or by cooling the steel sheet for about 60 to 180 seconds after the straining step (d) during which the steel sheet is maintained in essentially the same temperature range as the temperature range at the end of the quenching step (c);
(f) thereafter performing final cold-rolling with a rolling reduction of about 80 to 95 % to a predetermined final thickness; and (g) annealing the cold-rolled steel sheet for primary recrystallization and for decarburization, applying an annealing separation agent comprising MgO and effecting secondary-recrystallization annealing and purification-annealing.
(a) providing a hot-rolled steel sheet of a silicon steel having a composition containing about 0.01 to 0.15 % by weight of acid-soluble Al, about 0.005 to 0.04 % by weight of Sb, about 0.035 to 0.090 % by weight of C, about 2.5 to 4.5 %
by weight of Si, about 0.05 to 0.15 % by weight of Mn, 0.01 to 0.04 % by weight of S or Se and the balance being essentially Fe and inevitable impurities;
(b) softening-annealing the hot-rolled steel sheet at a temperature of about 850 to 1,200°C to dissolve AlN
before final cold rolling;
(c) successively quenching the softening-annealed steel sheet at a cooling speed of about 15 to 500°C/s to a temperature of about 500°C or less, thereby precipitating AlN
without precipitating C at grain boundaries;
(d) applying to the quenched steel sheet a strain ranging from about 0.005 to 3.0 % while maintaining the sheet at a temperature in the range from about the temperature reached by quenching to about 200°C;
(e) controlling carbide precipitation at an effective cooling speed of about 2°C/s or less by cooling the steel sheet during the straining step (d) or by cooling the steel sheet for about 60 to 180 seconds after the straining step (d) during which the steel sheet is maintained in essentially the same temperature range as the temperature range at the end of the quenching step (c);
(f) thereafter performing final cold-rolling with a rolling reduction of about 80 to 95 % to a predetermined final thickness; and (g) annealing the cold-rolled steel sheet for primary recrystallization and for decarburization, applying an annealing separation agent comprising MgO and effecting secondary-recrystallization annealing and purification-annealing.
2. A method according to claim 1, wherein the final sheet thickness is about 0.15 to 0.25 mm.
3. A method according to claim 1, wherein the final cold-rolling is performed at a temperature within the range of about 200 to 400°C.
4. A method according to claim 1, wherein the final cold-rolling step includes aging the steel sheet at a temperature in the range of about 200 to 400°C.
5. A method according to any one of claims 1 to 4, wherein the strain applying step is performed by applying a tension in the longitudinal direction of the steel that is in a lengthy sheet form.
6. A method according to any one of claims 1 to 4, wherein the strain applying step is performed by bending the steel sheet using a roll.
7. A method according to any one of claims 1 to 4, wherein the strain applying step is performed by applying shot blast.
8. A method according to any one of claims 1 to 4, wherein the quenching step (c) is performed to cool the softening-annealed steel sheet to a temperature in the range of from 450 to 300°C; and the strain is applied in step (d) in an amount of from 0.01 to 1.0 % in the longitudinal direction of the steel which is in a lengthy sheet form.
9. A method according to claim 8, wherein the strain applying step is performed by bending the steel sheet using a roll.
10. A method according to claim 8, wherein the strain applying step is performed by applying shot blast.
Applications Claiming Priority (2)
| Application Number | Priority Date | Filing Date | Title |
|---|---|---|---|
| JP2197822A JPH0784615B2 (en) | 1990-07-27 | 1990-07-27 | Method for producing grain-oriented silicon steel sheet with excellent magnetic flux density |
| JP197822/1990 | 1990-07-27 |
Publications (2)
| Publication Number | Publication Date |
|---|---|
| CA2048014A1 CA2048014A1 (en) | 1992-01-28 |
| CA2048014C true CA2048014C (en) | 1997-05-06 |
Family
ID=16380917
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| Application Number | Title | Priority Date | Filing Date |
|---|---|---|---|
| CA 2048014 Expired - Fee Related CA2048014C (en) | 1990-07-27 | 1991-07-26 | Method of manufacturing an oriented silicon steel sheet having improved magnetic flux density |
Country Status (5)
| Country | Link |
|---|---|
| EP (1) | EP0468819B1 (en) |
| JP (1) | JPH0784615B2 (en) |
| KR (1) | KR930011405B1 (en) |
| CA (1) | CA2048014C (en) |
| DE (1) | DE69109010T2 (en) |
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| JP3160281B2 (en) * | 1990-09-10 | 2001-04-25 | 川崎製鉄株式会社 | Method for producing grain-oriented silicon steel sheet with excellent magnetic properties |
| JP3275712B2 (en) * | 1995-10-06 | 2002-04-22 | 日本鋼管株式会社 | High silicon steel sheet excellent in workability and method for producing the same |
| JP5287615B2 (en) * | 2009-09-04 | 2013-09-11 | Ｊｆｅスチール株式会社 | Method for producing grain-oriented electrical steel sheet |
| CN102382963B (en) * | 2011-11-08 | 2013-11-27 | 北京科技大学 | Heat treatment method for improving room-temperature ductility of high-silicon electrical steel |
| CN104011246B (en) * | 2011-12-27 | 2016-08-24 | 杰富意钢铁株式会社 | Orientation electromagnetic steel plate |
| CN102787276B (en) * | 2012-08-30 | 2014-04-30 | 宝山钢铁股份有限公司 | High magnetic induction oriented silicon steel and manufacturing method thereof |
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| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| JPS5340982A (en) * | 1976-09-22 | 1978-04-13 | Toyo Kanetsu Kk | Automatic solid storage device |
| JPS58157917A (en) * | 1982-03-15 | 1983-09-20 | Kawasaki Steel Corp | Manufacture of unidirectional silicon steel plate with superior magnetic characteristic |
| JPS61124406A (en) * | 1984-11-21 | 1986-06-12 | Toshiba Corp | Picking device |
| DE3675945D1 (en) * | 1986-03-25 | 1991-01-10 | Kawasaki Steel Co | METHOD FOR PRODUCING THIN SILICON STEEL SHEETS WITH GOSS TEXTURE WITH LOW WATER LOSS AND EXCELLENT SURFACE PROPERTIES. |
| JPS63310406A (en) * | 1987-06-12 | 1988-12-19 | Seibu Electric & Mach Co Ltd | Method and device for cargo preparation of article |
| JP2561519B2 (en) * | 1988-09-09 | 1996-12-11 | 株式会社ダイフク | Warehouse equipment |
-
1990
- 1990-07-27 JP JP2197822A patent/JPH0784615B2/en not_active Expired - Fee Related
-
1991
- 1991-07-26 EP EP19910306880 patent/EP0468819B1/en not_active Expired - Lifetime
- 1991-07-26 DE DE1991609010 patent/DE69109010T2/en not_active Expired - Fee Related
- 1991-07-26 CA CA 2048014 patent/CA2048014C/en not_active Expired - Fee Related
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Also Published As
| Publication number | Publication date |
|---|---|
| EP0468819A1 (en) | 1992-01-29 |
| JPH0784615B2 (en) | 1995-09-13 |
| JPH0483823A (en) | 1992-03-17 |
| DE69109010D1 (en) | 1995-05-24 |
| KR930011405B1 (en) | 1993-12-06 |
| EP0468819B1 (en) | 1995-04-19 |
| CA2048014A1 (en) | 1992-01-28 |
| DE69109010T2 (en) | 1996-01-11 |
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