# Patent

## Title
Fe-based metal sheet and manufacturing method thereof

## Abstract

A cast slab containing C: less than 0.02 mass % and made of an Fe-based metal of an α-γ transforming component is subjected to hot rolling at a temperature of an A3 point or higher and is subjected to α-region rolling at a temperature of 300° C. or higher and lower than the A3 point, and thereby a base metal sheet having a {100} texture in a surface layer portion is fabricated. Then, by performing a heat treatment under predetermined conditions, an Fe-based metal sheet is obtained in which a Z value is not less than 2.0 nor more than 200 when intensity ratios of respective {001}<470>, {116}<6 12 1>, and {223}<692> directions in a sheet plane by X-ray diffraction are set to A, B, and C respectively and Z =(A+0.97B)/0.98C is satisfied.

## Background
TECHNICAL FIELD
The present invention relates to an Fe-based metal sheet having a high accumulation degree of {200} planes suitably used for magnetic cores and the like of electric motors, power generators, and transformers and capable of contributing to downsizing of these magnetic cores and reduction in energy loss, and a manufacturing method thereof. This application is based upon and claims the benefit of priority of the prior Japanese Patent Application No. 2011-100014, filed on Apr. 27, 2011, the prior Japanese Patent Application No. 2011-101893, filed on Apr. 28, 2011, and the prior Japanese Patent Application No. 2012-070166, filed on Mar. 26, 2012, the entire contents of which are incorporated herein by reference.
BACKGROUND ART
Electrical steel sheets alloyed with silicon or/and the like have been conventionally used for magnetic cores of electric motors, power generators, transformers, and the like. Among electrical steel sheets, non-oriented electrical steel sheets having relatively random crystal orientations can be manufactured at a low cost, to thus be used for motors, transformers, and the like of home electric appliances, and the like in a multipurpose manner. The crystal orientations of this non-oriented electrical steel sheet are random, thus making it impossible to obtain a high magnetic flux density. In contrast to this, grain-oriented electrical steel sheets having aligned crystal orientations can obtain a high magnetic flux density, to thus be applied to high-end use for driving motors and the like of HV vehicles and the like. However, in a manufacturing method of a grain-oriented electrical steel sheet that is industrialized currently, a long-time heat treatment is required, to thus increase the cost.
As above, in the non-oriented electrical steel sheet, a sufficiently high magnetic flux density cannot be obtained, and in the grain-oriented electrical steel sheet, the direction in which a high magnetic flux density can be obtained is limited to one to two direction/directions. On the other hand, in HV vehicles, and the like, achievement of high torque and downsizing are required, and there is a demand for manufacturing a metal sheet capable of obtaining a high magnetic flux density in an in-plane circumferential direction thoroughly as a metal sheet to be used for core materials of driving motors, and the like. Thus, as methods other than the industrialized manufacturing method of the grain-oriented electrical steel sheet, there have been proposed a technique of increasing an accumulation degree of a specific crystal orientation and various techniques of decreasing a core loss. However, in the technique described in Patent Literature 7, for example, it is possible to increase an accumulation degree of {200} planes, but directionality to a specific orientation occurs, to thus have a high magnetic flux density in a specific direction, but a high magnetic flux density cannot be obtained in an in-plane circumferential direction thoroughly, and the like, resulting in that in a conventional technique, satisfactory properties are not necessarily obtained.
CITATION LIST
Patent Literature
Patent Literature 1: Japanese Laid-open Patent Publication No. 10-168542
Patent Literature 2: Japanese Laid-open Patent Publication No. 2006-45613
Patent Literature 3: Japanese Laid-open Patent Publication No. 2006-144116
Patent Literature 4: Japanese Laid-open Patent Publication No. 10-180522
Patent Literature 5: Japanese Laid-open Patent Publication No. 01-252727
Patent Literature 6: Japanese Laid-open Patent Publication No. 07-173542
Patent Literature 7: International Publication Pamphlet No. WO2011/052654
SUMMARY OF INVENTION
Technical Problem
Thus, an object of the present invention is to provide an Fe-based metal sheet that is likely to become magnetized in a sheet plane and further has a texture capable of obtaining a high magnetic flux density thoroughly in an in-plane circumferential direction, and a manufacturing method thereof.
Solution to Problem
The present inventors, as a result of earnest examination, found that an orientation ratio to a specific orientation is controlled with respect to an Fe-based metal of an iron sheet or the like, and thereby a <100> orientation in αFe is more densely and thoroughly distributed in a metal sheet plane to make it possible to obtain a high magnetic flux density thoroughly in an in-plane circumferential direction.
Further, the present inventors conceived that in order to manufacture such an Fe-based metal sheet, a texture in which an accumulation degree of {100} planes is increased is first formed in a surface layer portion, and at the time of γ-α transformation by the subsequent heat treatment, the texture is transformed while taking over its {100} texture. Then, they earnestly examined a method of forming the {100} texture in the surface layer portion and achievement of high accumulation of {200} planes using the γ-α transformation.
As a result, it was found that when the Fe-based metal sheet is manufactured from a slab by rolling, a rolling temperature and a reduction ratio are optimized, thereby making it possible to form the {100} texture in at least the surface layer portion. Then, it was found that when the {100} texture in the surface layer portion is taken over by using the γ-αtransformation thereafter, a different metal except Fe is made to diffuse beforehand from the surface and a diffused region is turned into an α-Fe phase, and thereby in the region turned into the α-Fe phase, the {100} texture is formed, and at the time of the γ-α transformation, an accumulation degree of {200} planes in the α-Fe phase further generated by the transformation increases and the <100> orientation is distributed more densely and thoroughly, thereby making it possible to obtain a high magnetic flux density in the in-plane circumferential direction thoroughly.
Further, the present inventors found that in the case of a large amount of C content being contained, when the C content is decreased by decarburization annealing, the decarburization annealing is performed under predetermined conditions, thereby also making it possible to form the {100} texture in at least the surface layer portion, and in the Fe-based metal sheet obtained finally, the <100> orientation is distributed more densely and thoroughly, thereby making it possible to obtain a high magnetic flux density in the in-plane circumferential direction thoroughly.
The gist of the present invention made as a result of such examinations is as follows.
(1) An Fe-based metal sheet, includes: at least one type of ferrite-forming element except Fe, in which when intensity ratios of respective {001}<470>, {116}<6 12 1>, and {223}<692> directions in a sheet plane by X-ray diffraction are set to A, B and C respectively and Z=(A+0.97B)/0.98C is satisfied, a Z value is not less than 2.0 nor more than 200.
(2) The Fe-based metal sheet according to (1), in which the ferrite-forming element diffuses from a surface to be alloyed with Fe.
(3) The Fe-based metal sheet according to (1) or (2), in which a layer containing the ferrite-forming element is formed on at least one side of surfaces of the Fe-based metal sheet, and the ferrite-forming element that has diffused from part of the layer is alloyed with Fe.
(4) The Fe-based metal sheet according to (3), in which a thickness of the layer containing the ferrite-forming element is not less than 0.01 μm nor more than 500 μm.
(5) The Fe-based metal sheet according to any one of (1) to (4), in which an accumulation degree of {200} planes is not less than 30% nor more than 99%, and an accumulation degree of {222} planes is not less than 0.01% nor more than 30%.
(6) The Fe-based metal sheet according to any one of (1) to (5), in which the ferrite-forming element is one type of element or more selected from a group consisting of Al, Cr, Ga, Ge, Mo, Sb, Si, Sn, Ta, Ti, V, W, and Zn.
(7) The Fe-based metal sheet according to any one of (1) to (6), in which at least a partial region including the surfaces of the Fe-based metal sheet is an α single phase region made of an α single phase based component, and a ratio of the α single phase region to a cross section of the Fe-based metal sheet is 1% or more.
(8) The Fe-based metal sheet according to any one of (1) to (7), in which a thickness of the Fe-based metal sheet is not less than 10 μm nor more than 6 mm.
(9) The Fe-based metal sheet according to any one of (1) to (8), in which the α single phase region is formed on a front surface side and a rear surface side of the Fe-based metal sheet, and a crystal grain straddling the α single phase region on the front surface side and the α single phase region on the rear surface side is formed.
(10) A manufacturing method of an Fe-based metal sheet, includes:
performing hot rolling on a cast slab containing C: less than 0.02 mass % and made of an Fe-based metal of an α-γ transforming component at a temperature of an A3 point of the cast slab or higher to obtain a hot-rolled sheet;
performing α-region rolling on the hot-rolled sheet at a temperature of higher than 300° C. and lower than the A3 point of the cast slab to obtain a rolled sheet;
performing cold rolling on the rolled sheet to obtain a base metal sheet having a thickness of not less than 10 μm nor more than 6 mm;
bonding a ferrite-forming element to one surface or both surfaces of the base metal sheet;
heating the base metal sheet having had the ferrite-forming element bonded thereto up to an A3 point of the base metal sheet; and
further heating the heated base metal sheet to a temperature of not lower than the A3 point of the base metal sheet nor higher than 1300° C. and holding the base metal sheet; and
cooling the heated and held base metal sheet to a temperature of lower than the A3 point of the base metal sheet.
(11) The manufacturing method of the Fe-based metal sheet according to (10), in which a reduction ratio in the α-region rolling is −1.0 or less in terms of true strain, and the sum of the reduction ratio in the α-region rolling and a reduction ratio in the cold rolling is −2.5 or less in terms of true strain.
(12) The manufacturing method of the Fe-based metal sheet according to (10) or (11), in which
a reduction ratio in the hot rolling is −0.5 or less in terms of true strain.
(13) A manufacturing method of an Fe-based metal sheet, includes:
heating a steel sheet containing C: not less than 0.02 mass % nor more than 1.0 mass %, having a thickness of not less than 10 μm nor more than 6 mm, and made of an Fe-based metal of an α-γ transforming component to a temperature of an A1 point or higher and a temperature at which a structure is turned into an α single phase when decarburization is performed until C becomes less than 0.02 mass %, to obtain a base metal sheet that has been subjected to decarburization in a range of not less than 5 μm nor more than 50 μm in a depth direction from its surface until C becomes less than 0.02 mass %;
bonding a ferrite-forming element to one surface or both surfaces of the base metal sheet;
heating the base metal sheet having had the ferrite-forming element bonded thereto up to an A3 point of the base metal sheet; and
further heating the heated base metal sheet to a temperature of not lower than the A3 point of the base metal sheet nor higher than 1300° C. and holding the base metal sheet; and
cooling the heated and held base metal sheet to a temperature of lower than the A3 point of the base metal sheet.
(14) The manufacturing method of the Fe-based metal sheet according to (13), in which the steel sheet made of the Fe-based metal further contains Mn of 0.2 mass % to 2.0 mass %, and decarburization and demanganization are performed in a combined manner.
(15) The manufacturing method of the Fe-based metal sheet according to (13) or (14), further includes:
performing carburization on a steel sheet containing C: less than 0.02 mass %, having a sheet thickness of not less than 10 μm nor more than 6 mm, and made of an Fe-based metal of an α-γ transforming component to control C to not less than 0.02 mass % nor more than 1.0 mass %.
Advantageous Effects of Invention
According to the present invention, it is possible to manufacture an Fe-based metal sheet capable of obtaining a high magnetic flux density thoroughly in an in-plane circumferential direction.



BRIEF DESCRIPTION OF DRAWINGS
FIG. 1 is a view for explaining a method of calculating an average magnetic flux density B50;
FIG. 2 is a conceptual diagram showing the relationship between a Z value and a ratio B50/Bs of the average magnetic flux density B50 to a saturation magnetic flux density Bs and a magnetic flux density difference ΔB;
FIG. 3A is a view schematically showing a structure of a cross section of a base metal sheet having a {100} texture formed in a surface layer portion;
FIG. 3B is a view schematically showing the structure of the cross section of the base metal sheet having a different metal layer formed in the surface layer portion;
FIG. 3C is a view schematically showing the structure of the cross section of the base metal sheet in a temperature increasing process;
FIG. 3D is a view schematically showing the structure of the cross section of the base metal sheet in a heating and holding process;
FIG. 3E is a view schematically showing the structure of the cross section of the base metal sheet in a cooling process;
FIG. 4A is a view schematically showing the structure of the cross section of the base metal sheet in a state of being held at a temperature of an A3 point or higher;
FIG. 4B is a view schematically showing the structure of the cross section of the base metal sheet after cooling in the case when the different metal layers are made to remain;
FIG. 4C is a view schematically showing the structure of the cross section of the base metal sheet in the case when the base metal sheet is alloyed up to its center portion in a state of being held at the temperature of the A3 point or higher;
FIG. 4D is a view schematically showing the structure of the cross section of the base metal sheet after cooling in the case when the base metal sheet is alloyed up to the center portion; and
FIG. 5 is a view schematically showing the structure of the cross section of the base metal sheet in which a crystal grain becomes coarse.



DESCRIPTION OF EMBODIMENTS
Generally, an orientation of easy magnetization exists in α-Fe crystal, and when in a direction in which direction cosines between <100>, <010>, <001> orientations, (which will be called a [100] orientation generically), and the orientation are large, excitation is performed in a fixed magnetic field and magnetometry is performed, a high magnetic flux density is likely to be obtained. On the other hand, when in a direction in which direction cosines with respect to a <111> orientation being an orientation of hard magnetization are large, excitation is performed and magnetometry is performed, a high magnetic flux density is unlikely to be obtained. The present inventors found that more [100] orientations in the α-Fe crystal exist in a sheet plane and further the α-Fe crystal is controlled to a specific texture that is thoroughly distributed in the sheet plane, and thereby direction cosines with respect to the [100] orientation always become large in an arbitrary direction in the metal sheet plane, and when a magnetic field is applied in an arbitrary direction in the metal sheet plane and magnetometry is performed, a high magnetic flux density can be obtained.
It is characterized in that a specific texture that an Fe-based metal sheet of the present invention has contains at least one type of ferrite-forming element except Fe, in which when intensity ratios in respective {001}<470>, {116}<6 12 1>, and {223}<692> directions in a sheet plane by X-ray diffraction are set to A, B, and C respectively and Z=(A+0.97B)/0.98C is satisfied, a Z value is not less than 2.0 nor more than 200.
Next, the previously described Z value will be explained.
The main orientations on which attention is focused in the present invention are {001}<470>, {116}<6 12 1>, and {223}<692>. When examining the state of a three-dimensional texture calculated by a vector method, the present inventors noticed that X-ray random intensity ratios in the above-described three plane orientations change depending on a magnetic property of a product, and learned that mathematizing this makes it possible to quantify the relationship with a magnetic property of a product and reached the present invention.
The X-ray random intensity ratios of these respective orientations may be obtained from a three-dimensional texture calculated by a vector method based on a pole figure of {110}, or may also be obtained from a three-dimensional texture calculated by a series expansion method using a plurality (preferably three or more) of pole figures out of pole figures of {110}, {100}, {211}, and {310}. For the X-ray random intensity ratios in the above-described respective crystal orientations by the latter method, for example, intensities of (001)[4 −7 0], (116)[1 −12 1], and (223)[6 −9 2] at a φ2=45° cross-section of the three-dimensional texture may be used as they are.
Subsequently, there will be explained a reason for which the expression of Z=(A+0.97B)/0.98C was found.
First, the intensity of the {001}<470> orientation is set to A. This orientation is in the {100} plane, so that direction cosines with respect to the {100} plane are 1.0. In the {100} plane, the [100] orientation being the orientation of easy magnetization exists, and thus orientation of this plane in the metal sheet plane is advantageous for obtaining a high magnetic flux density in the metal sheet plane. Thus, the intensity A is weighted with the direction cosines of 1.0 in terms of the degree of contribution to improving a magnetic flux density to be set to one of parameters in the Z value.
Next, the intensity of the {116}<6 12 1> orientation is set to B. An angular difference between this orientation and the {001} plane is 13.3° and direction cosines are 0.97. In the {001} plane as well, the [100] orientation being the orientation of easy magnetization exists, and thus orientation of this plane in the metal sheet plane is advantageous for obtaining a high magnetic flux density in the metal sheet plane. For this reason, the intensity B is weighted with the direction cosines of 0.97 in terms of the degree of contribution to improving a magnetic flux density to be set to one of parameters in the Z value.
Further, the intensity of the {223}<692> orientation is set to C. An angular difference between the {223}<692> orientation and a {111} plane is 11.4° and direction cosines are 0.98. As described previously, in the {111} plane, the [100] orientation being the orientation of easy magnetization is not contained, and orientation of this plane in the metal sheet plane is disadvantageous for obtaining a high magnetic flux density. Thus, the intensity C is set not to have the degree of contribution to improving a magnetic flux density, is put in the Z value as a parameter that performs division, and is multiplied by 0.98 being the direction cosines with respect to the {111} plane as its weighting.
From the above thought, it was found that when the intensity ratios in the respective {001}<470>, {116}<6 12 1>, and {223}<692> directions in the metal sheet plane by X-ray diffraction are set to A, B, and C respectively, the expression of Z=(A+0.978)/0.98C is created, and as the Z value is increased, a high magnetic flux density can be obtained when excitation is performed in the metal sheet plane to perform magnetometry.
Further, the present inventors were able to find from a large number of experiments that a special condition capable of obtaining a high magnetic flux density in an arbitrary direction in the metal sheet plane is that the Z value is not less than 2.0 nor more than 200. They grasped the fact that the Z value is limited to this range, and thereby the [100] orientation being the orientation of easy magnetization is thoroughly distributed in the metal sheet plane, but have not obtained evidence making theoretical explanation of this phenomenon possible so far.
The present inventors found that when the Z value is not less than 2.0 nor more than 200, a ratio B50/Bs of an average magnetic flux density B50 to a saturation magnetic flux density Bs becomes a high level of 0.80 or more and a magnetic flux density difference ΔB measured in the metal sheet plane becomes a low level of 0.15 T or less. FIG. 2 schematically shows this relationship.
When the Z value is less than 2.0, crystal orientation of α-Fe shows a tendency to decrease the [100] orientations being the orientation of easy magnetization in the metal sheet plane. Alternately, it shows a tendency that the distribution of the [100] orientations in the metal sheet plane becomes non-uniform. That is, the average magnetic flux density B50 in the metal sheet plane becomes small and the ratio B50/Bs of the average magnetic flux density B50 to the saturation magnetic flux density Bs becomes less than 0.8. Alternately, only the magnetic flux density in a specific direction increases and the magnetic flux density difference ΔB becomes greater than 0.15 T. Thus, the Z value is set to 2.0 or more in the present invention.
On the other hand, when the Z value exceeds 200, the increase in the magnetic flux density is saturated and an increase in uniformity of the magnetic flux density in the metal sheet plane is also saturated. In contrast to this, in order to manufacture a metal sheet such that the Z value exceeds 200, a heat treatment time is prolonged, or the like, which becomes difficult industrially, and thus the condition of the Z value is set to 200 or less.
Here, FIG. 1 is a view for explaining a method of calculating the average magnetic flux density B50. A manufacturing method will be described later, but it is found that α-region rolling is performed at 800° C. and as a different metal, 2.6 mass % of Sn and 0.9 mass % of Al are used, and thereby in an obtainable Fe-based metal sheet having a thickness of 0.2 mm, a high magnetic flux density can be obtained thoroughly in an in-plane circumferential direction.
Here, in a metal sheet having a higher accumulation degree of {200} planes among textures of the Fe-based metal sheet of the present invention in which the Z value is not less than 2.0 nor more than 200, a higher magnetic flux density can be obtained. Specifically, in a texture in which an accumulation degree of {200} planes in an α-Fe phase is not less than 30% nor more than 99% and an accumulation degree of {222} planes in the α-Fe phase is not less than 001% nor more than 30%, a higher magnetic flux density can be obtained.
When the accumulation degree of the {200} planes is less than 30% or the accumulation degree of the {222} planes is greater than 30%, the average magnetic flux density B50 tends to slightly decrease even though the Z value is in the present invention range. Further, in a metal sheet in which the accumulation degree of the {200} planes is greater than 99% or the accumulation degree of the {222} planes is less than 0.01%, the increase in the magnetic flux density B50 is saturated and a heat treatment time is prolonged, and the like, resulting in that manufacturing conditions become disadvantageous industrially.
Next, the manufacturing method of the previously described Fe-based metal sheet will be explained.
(First Embodiment)
As a manufacturing method of an Fe-based metal sheet in this embodiment, a rolling temperature and a reduction ratio are optimized, and thereby a {100} texture is formed in at least a surface layer portion of the metal sheet, a ferrite-forming element is made to diffuse into this partial or whole region from its surface, and at the time of cooling, the whole Fe-based metal sheet is oriented in {100}. This makes it possible to obtain a high magnetic flux density in an arbitrary direction in a metal sheet plane.
This embodiment as above is based on the fact found by the present inventors that {100} crystal grains in the texture formed in the surface layer portion preferentially grow at an A3 point or higher in a heating process to be performed for the diffusion of the ferrite-forming element, and further when the ferrite-forming element is made to diffuse into the inner portion to make the Fe-based metal sheet alloyed therewith and then cooling is performed, an accumulation degree of {200} planes in the sheet plane of the Fe-based metal sheet increases.
[Explanation of the Basic Principle of the First Embodiment of the Present Invention]
First, the basic principle of this embodiment capable of obtaining a high accumulation degree of {200} planes will be explained based on FIG. 3A to FIG. 3E.
(a) Manufacture of a Base Metal Sheet (Seeding of a Texture)
In a process in which a cast slab containing C: less than 0.02 mass % and made of an Fe-based metal of an α-γ transforming component is decreased in thickness by rolling and thereby a metal sheet is obtained, hot rolling is performed at a sheet temperature of the A3 point or higher, α-region rolling is performed at a sheet temperature of lower than the A3 point and 300° C. or higher, and further cold rolling is performed to a predetermined sheet thickness. By this process, as shown in FIG. 3A, a base metal sheet 1 having an inner region 4 made of Fe in an α phase and having a {100} texture 2 in at least a surface layer portion 3 is obtained. Further, a seed of crystal that satisfies the condition of the Z value is formed in a recrystallized texture by a particular deformation slip.
(b) (Formation of a Second Layer)
Next, as shown in FIG. 3B, the ferrite-forming element such as Al, for example, is bonded to one surface or both surfaces of the cold-rolled base metal sheet 1 by using a vapor deposition method or the like to form a second layer 5.
(c) Saving of the Texture
Next, the base metal sheet 1 having had the ferrite-forming element bonded thereto is heated to the A3 point of the base metal sheet 1 to make the ferrite-forming element diffuse into the partial or whole region having the {100} texture 2 in the base metal sheet 1, to make the base metal sheet 1 alloyed therewith. As shown in FIG. 3C, an alloyed region 6 is transformed to the α phase from a γ phase to have an α single phase component. At this time, the alloyed region 6 is transformed while taking over orientation of the {100} texture 2 formed in the surface layer portion 3, so that a structure oriented in {100} is formed also in the alloyed region 6.
(d) Achievement of High Accumulation of the Texture
Next, the partially alloyed base metal sheet 1 is further heated to a temperature of not lower than the A3 point nor higher than 1300° C. and the temperature is held. The region of the α single phase component is an α-Fe phase not undergoing γ transformation, and thus the {100} crystal grains are maintained as they are, the {100} crystal grains preferentially grow in the region, and the accumulation degree of the {200} planes increases. Further, as shown in FIG. 3D, a region 8 not having the α single phase component is transformed to the γ phase from the α phase.
Further, when a holding time of the temperature after the heating is prolonged, the {100} crystal grains are united to preferentially grow to large {100} crystal grains 7. As a result, the accumulation degree of the {200} planes further increases. Further, with the diffusion of the ferrite-forming element, the region 6 alloyed with the ferrite-forming element is transformed to the α phase from the γ phase. At this time, in the region adjacent to the region to be transformed, crystal grains in the α phase oriented in {100} are already formed, and at the time of the transformation to the α phase from the γ phase, the region 6 is transformed while taking over a crystal orientation of the adjacent crystal grains in the α phase. Thereby, the holding time is prolonged and the accumulation degree of the {200} planes increases.
(e) Growth of the Texture
The base metal sheet is cooled to a temperature of lower than the A3 point. At this time, as shown in FIG. 3E, a γ-Fe phase in an unalloyed inner region 10 is transformed to the α-Fe phase. This inner region 10 is adjacent to the region in which the crystal grains in the α phase oriented in {100} are already formed in a temperature region of the A3 point or higher, and at the time of the transformation to the α phase from the γ phase, the inner region 10 is transformed while taking over the crystal orientation of the adjacent crystal grains in the α phase and larger crystal grains 9 in the α phase oriented in {100} are formed. Therefore, the accumulation degree of the {200} planes increases also in the region. By this phenomenon, the high accumulation degree of the {200} planes can be obtained even in the unalloyed region.
When at the stage of the preceding state shown in FIG. 3D, the temperature of the A3 point or higher is held until the whole metal sheet is alloyed, the structure having the high accumulation degree of the {200} planes is already formed in the whole metal sheet, and thus the cooling is performed while the state when the cooling is started is maintained.
In the above, the basic principle of this embodiment was explained, and there will be further explained a limiting reason of each condition that defines the manufacturing method of this embodiment and preferable conditions of this embodiment.
[Fe-Based Metal to be the Base Material] (C Content)
In this embodiment, first, crystal grains oriented in {100} to serve as seeds for increasing the accumulation degree of the {200} planes in the sheet are formed in the surface layer portion of the base metal sheet made of the Fe-based metal. Then, the γ-α transformation is made to progress in the metal sheet while taking over a crystal orientation of the crystal grains in the α phase to serve as the seeds finally, to thereby increase the accumulation degree of the {200} planes of the whole metal sheet. For this reason, the Fe-based metal used for the base metal sheet has a composition of the α-γtransforming component. When the Fe-based metal used for the base metal sheet has the α-γ transforming component, the ferrite-forming element is made to diffuse into the metal sheet to make the metal sheet alloyed therewith, thereby making it possible to form the region having the α single phase based component.
In this embodiment, the C content of the base metal sheet is set to less than 0.02 mass %. Further, in terms of a magnetic property of a product metal sheet, the C content is preferably 0.01 mass % or less. Under the condition of the C content being less than 0.02 mass %, the ferrite-forming element is made to diffuse into the metal sheet to make the metal sheet alloyed therewith, thereby making it possible to form the region having the α single phase based component. Incidentally, C is a component to remain in a process of manufacturing the slab and the less C is, the more preferred it is in terms of the magnetic property, and thus its lower limit is not necessary needed, but it is preferably set to 0.0001 mass % or more in terms of the cost of a refining process.
(Other Containing Elements)
In principle, being applicable to the Fe-based metal having the α-γ transforming component, this embodiment is not limited to the Fe-based metal in a specific composition range. Typical examples of the α-γ transforming component are pure iron, steel such as ordinary steel, and the like. For example, it is a component containing pure iron or steel containing C of 1 ppm to less than 0.02 mass % as described above and a balance being composed of Fe and inevitable impurities as its base and containing an additive element as required. Instead, it may be silicon steel of the α-γ transforming component having C: less than 0.02 mass % and Si: 0.1 mass % to 2.5 mass % as its basic component. Further, as other impurities, a trace amount of Ni, Cr, Al, Mo, W, V, Ti, Nb, B, Cu, Zr, Y, Hf, La, Ce, N, O, P, S, and/or the like are/is contained. Further, Al and Mn are added to increase electric resistance, to thereby decrease a core loss, and Co is added to increase the saturation magnetic flux density Bs, to thereby increase a magnetic flux density, which are also included in the present invention range.
(Thickness of the Base Metal Sheet)
The thickness of the base metal sheet is set to not less than 10 μm nor more than 6 mm. When the thickness is less than 10 μm, when the base metal sheets are stacked to be used as a magnetic core, the number of the sheets to be staked is increased to increase gaps, resulting in that a high magnetic flux density cannot be obtained. Further, when the thickness exceeds 6 mm, it is not possible to make the {100} texture grow sufficiently even though a reduction ratio of the α-region rolling is adjusted, resulting in that a high magnetic flux density cannot be obtained.
[Rolling Conditions]
In this embodiment, as described previously, the Fe-based metal having, in at least the surface layer portion, the crystal grains oriented in {100} to serve as the seeds for increasing the accumulation degree of the {200} planes in the metal sheet is used as a starting material. As a method of achieving high accumulation of the {100} planes of the base metal sheet, a method of performing α-region rolling in a process in which a cast slab is rolled to a sheet shape is used.
First, a cast slab containing C: less than 0.02 mass % and made of the Fe-based metal of the α-γtransforming component such as a continuous cast slab or an ingot is prepared. Then, in a process in which the cast slab is decreased in thickness by rolling to obtain the base metal sheet, first the hot rolling is performed at a temperature of the A3 point or higher. Next, the α-region rolling is performed at a temperature of lower than the A3 point and higher than 300° C., and further the metal sheet is subjected to cold rolling to a predetermined thickness, and thereby the base metal sheet having the {100} texture formed in the surface layer portion is obtained.
As for a reduction ratio in each of rolling processes to be performed until the base metal sheet is obtained from the cast slab, the total reduction ratio in the α-region rolling is preferably set to −1.0 or less in terms of true strain and the sum of the total reduction ratio in the α-region rolling and the total reduction ratio in the cold rolling is preferably set to −2.5 or less in terms of true strain. Conditions other than these may create a possibility that the {100} texture cannot be sufficiently formed in the surface layer portion. A method of expressing the reduction ratio by true strain E is expressed by the following expression (1), where in each of the rolling processes, the thickness before the rolling is set to h0 and the thickness after the rolling is set to h.

ε=ln(h/h0)  (1)

When the sum of the total reduction ratio in the α-region rolling and the total reduction ratio in the cold rolling is in the previously described preferred range, a deformed structure in which the {100} texture is formed by recrystallization can be provided to at least the vicinity of the surface layer portion of the base metal sheet. Particular crystal slip and crystal rotation to occur at these reduction ratios are thought to occur. Thus, they are preferably in these ranges.
Further, as for the reduction ratio in each of the rolling processes to be performed until the base metal sheet is obtained from the cast slab, the reduction ratio in the hot rolling is preferably −0.5 or less in terms of true strain, thereby making it easier to obtain the higher accumulation degree of the {200} planes. This results from the fact found by the present inventors that in order that desirable deformation should be performed in the α-region rolling and the cold rolling, deformation in the hot rolling in a γ region is also closely affected. Thus, these ranges are preferred.
The region of the surface layer portion in which the {100} texture is formed preferably has 1 μm or more of a distance in a sheet thickness direction from the surface. Thereby, it is possible to bring the accumulation degree of the {200} planes to 30% or more in the following diffusion treatment. The upper limit of the distance is not limited in particular, but it is difficult to form the {100} texture in a region of 500 am or more by rolling.
Incidentally, the measurement of the accumulation degree of the {200} planes can be performed by X-ray diffraction using a MoKα ray. To be in more detail, in the α-Fe crystal, integrated intensities of 11 orientation planes ({110}, {200}, {211}, {310}, {222}, {321}, {411}, {420}, {332}, {521}, and {442}) parallel to a sample surface are measured for each sample, each measured value is divided by a theoretical integrated intensity of the sample having a random orientation, and thereafter, a ratio of the intensity of {200} or {222} is obtained in percentage.
At this time, for example, the accumulation degree of the {200} planes is expressed by Expression (2) below.

accumulation degree of {200} planes=[{i(200)/I(200)}/Σ{i(hkl)/I(hkl)}]×100  (2)

Here, i(hkl) is an actually measured integrated intensity of {hkl} planes in the measured sample, and I(hkl) is a theoretical integrated intensity of the {hkl} planes in the sample having the random orientation. Further, Σ is the sum of the 11 orientation planes in the α-Fe crystal. Here, instead of the theoretical integrated intensity of the sample having the random orientation, actually measured values using the sample may be used.
[Different Metal]
Next, a different metal except Fe is made to diffuse into the base metal sheet manufactured by the above-described rolling processes to increase the region of the {100} texture in the thickness direction of the steel sheet. As the different metal, the ferrite-forming element is used. As a procedure, first, the different metal is bonded in a layered form as the second layer to one surface or both surfaces of the base metal sheet made of the Fe-based metal of the α-γ transforming component. Then, a region alloyed by having had elements of the different metal diffuse thereinto is turned to have the α single phase based component and to be able to be maintained as not only the region transformed to the α phase, but also a seed oriented in {100} for increasing the accumulation degree of the {200} planes in the metal sheet. As such a ferrite-forming element, at least one type of Al, Cr, Ga, Ge, Mo, Sb, Si, Sn, Ta, Ti, V, W, and Zn can be used alone or in a combined manner.
As a method of bonding the different metal in a layered form to the surface of the base metal sheet, there can be employed various methods such as a plating method of hot dipping, electrolytic plating, or the like, a rolling clad method, a dry process of PVD, CVD, or the like, and further powder coating. As a method of efficiently bonding the different metal for industrially implementing the method, the plating method or the rolling clad method is suitable.
The thickness of the different metal before the heating when the different metal is bonded is preferably not less than 0.05 μm nor more than 1000 μm. When the thickness is less than 0.05 μm, it is not possible to obtain the sufficient accumulation degree of the {200} planes. Further, when the thickness exceeds 1000 μm, even when the different metal layer is made to remain, its thickness becomes larger than necessary.
[Heating and Diffusion Treatment]
The base metal sheet having had the ferrite-forming element as the different metal bonded thereto is heated up to the A3 point of the base metal sheet, to thereby make the ferrite-forming element diffuse into the partial or whole region of the {100} texture formed in the surface layer portion of the base metal sheet to make the base metal sheet alloyed therewith. The region alloyed with the ferrite-forming element is turned to have the α single phase component and the region is transformed to the α phase from the γ phase. At this time, the region is transformed while taking over the orientation of the {100} texture formed in the surface layer portion, and thus the structure oriented in {100} is formed also in the alloyed region. As a result, in the alloyed region, a structure in which the accumulation degree of the {200} planes in the α-Fe phase becomes not less than 25% nor more than 50% and in accordance with it, the accumulation degree of the {222} planes in the α-Fe phase becomes not less than 1% nor more than 40% is formed.
Then, the base metal sheet is further heated to a temperature of not lower than the A3 point nor higher than 1300° C. and the temperature is held. The region alloyed already is turned into an α single phase structure that is not transformed to the γphase, so that the crystal grains in the {100} texture are maintained as they are, and in the region, the crystal grains in the {100} texture preferentially grow and the accumulation degree of the {200} planes increases. Further, the region not having the α single phase component is transformed to the γ phase.
Further, when the holding time is prolonged, the crystal grains in the {100} texture are united to one another to preferentially grow. As a result, the accumulation degree of the {200} planes further increases. Further, with the further diffusion of the ferrite-forming element, the region alloyed with the ferrite-forming element is transformed to the α phase from the γ phase. At this time, as shown in FIG. 4A, in the regions adjacent to the regions to be transformed, crystal grains 7 in the α phase oriented in {100} are already formed, and at the time of the transformation to the α phase from the γ phase, the regions alloyed with the ferrite-forming element are transformed while taking over a crystal orientation of the adjacent crystal grains 7 in the α phase. By these phenomena, the holding time is prolonged and the accumulation degree of the {200} planes increases. Further, as a result, the accumulation degree of the {200} planes decreases.
Incidentally, in order to finally obtain the high accumulation degree of the {200} planes of 50% or more, it is preferred that the holding time should be adjusted to, at this stage, bring the accumulation degree of the {200} planes in the α-Fe phase to 30% or more and bring the accumulation degree of the {200} planes in the α-Fe phase to 30% or less. Further, when the A3 point or higher is held until the whole metal sheet is alloyed, as shown in FIG. 4C, the α single phase structures are formed up to the center portion of the metal sheet and grain structures oriented in {100} reach the center of the metal sheet.
A holding temperature after the temperature is increased is set to not lower than the A3 point nor higher than 1300° C. Even when the metal sheet is heated at a temperature higher than 1300° C., an effect with respect to the magnetic property is saturated. Further, cooling may be started immediately after the temperature reaches the holding temperature (in the case, the temperature is held for 0.01 second or longer substantially), or cooling may also be started after the temperature is held for 600 minutes or shorter. Even when the temperature is held for longer than 600 minutes, the effect is saturated. When this condition is satisfied, the achievement of high accumulation of the seeds oriented in the {200} plane further progresses to make it possible to more securely bring the accumulation degree of the {200} planes in the α-Fe phase to 30% or more after the cooling.
[Cooling after the Heating and Diffusion Treatment]
After the diffusion treatment, when the cooling is performed while the region that is not alloyed with the ferrite-forming element is remaining, as shown in FIG. 4B, at the time of the transformation to the α phase from the γ phase, the unalloyed region is transformed while taking over the crystal orientation of the regions in which the crystal grains 9 in the α phase oriented in {100} are already formed. Thereby, the accumulation degree of the {200} planes increases, and the metal sheet having the texture in which the accumulation degree of the {200} planes in the α-Fe phase is not less than 30% nor more than 99% and the accumulation degree of the {200} planes in the α-Fe phase is not less than 0.01% nor more than 30% is obtained, the crystal satisfying the condition of the Z value grows, and a high magnetic flux density can be obtained in an arbitrary direction in the metal sheet plane.
Further, as shown in FIG. 4C, when the A3 point or higher is held until the whole metal sheet is alloyed, and the grain structures oriented in {100} reach the center of the metal sheet, as shown in FIG. 4D, the metal sheet is cooled as it is, and the texture in which the crystal grains 9 oriented in {100} reach up to the center of the metal sheet can be obtained. Thereby, the whole metal sheet is alloyed with the different metal, and the metal sheet having the texture in which the accumulation degree of the {200} planes in the α-Fe phase is not less than 30% nor more than 99% and the accumulation degree of the {222} planes in the α-Fe phase is not less than 0.01% nor more than 30% is obtained.
As above, the value of the accumulation degree of the {200} planes and the remaining state of the different metal on the surface of the base metal sheet change depending on the holding time of the temperature of the A3 point or higher and the holding temperature. The example shown in FIG. 4B is in a state where the grain structures oriented in {100} do not reach up to the center of the metal sheet, the different metal also remain on the surfaces, and an α single phase front surface side region and an α single phase rear surface side region being the second layer are formed, but it is also possible to obtain the grain structures oriented in {100} up to the center of the metal sheet and to alloy all the second layers on the surfaces.
Incidentally, at the time of the cooling after the diffusion treatment, a cooling rate is preferably not less than 0.1° C./sec nor more than 500° C./sec. When the cooling is performed in this temperature range, the growth of the seeds oriented in the {200} plane further progresses.
Further, when the second layers are made to remain on the obtainable Fe-based metal sheet having a thickness of not less than 10 μm nor more than 6 mm, the thickness of the second layer is preferably set to not less than 0.01 μm nor more than 500 μm. Further, a ratio of the α single phase region alloyed at this stage is preferably 1% or more in a cross section of the Fe-based metal sheet.
Further, at the time of cooling to a temperature of lower than the A3 point in the state shown in FIG. 4A, an average cooling rate at the time of cooling to the A3 point—50° C. from the A3 point may be set to 50° C./minute or less. When the cooling is performed at the cooling rate in this range, the adjacent crystal grains oriented in {100} are united to one another to grow, and as shown in FIG. 5, a coarse crystal grain 11 straddling part of an α single phase front surface side region 6a adjacent to a front surface side second layer 5a and straddling part of an α single phase rear surface side region 6b adjacent to a rear surface side second layer 5b is formed. When the average cooling rate from the A3 point to the A3 point—50° C. becomes greater than 50° C./minute, there is no sufficient time for growth of the crystal grain 11, resulting in that an excellent core loss property cannot be obtained. On the other hand, the lower limit of the average cooling rate from the A3 point to the A3 point—50° C. is not limited, but the lower limit is preferably set to 1° C./minute in terms of the productivity.
Further, in order to obtain a more excellent core loss property, an average cooling rate at the time of cooling to the A3 point—10° C. from the A3 point is preferably set to 20° C./minute or less. On the other hand, the lower limit of the average cooling rate from the A3 point to the A3 point—10° C. is not limited, but the lower limit is preferably set to 1° C./minute in terms of the productivity.
(Second Embodiment)
In the previously described first embodiment, there was explained the manufacturing method of the previously described Fe-based metal sheet by using the cast slab containing C: less than 0.02 mass % and made of the Fe-based metal of the α-γ transforming component. In contrast to this, in this embodiment, there will be explained a manufacturing method of the previously described Fe-based metal sheet by using a cast slab containing C: 0.02 mass % or more.
When the C content is large, a good magnetic property cannot be obtained, so that it is necessary to remove C by performing decarburization annealing. Thus, the decarburization annealing is performed under conditions to be explained below, thereby making it possible to increase the accumulation degree of the {200} planes.
In the method of this embodiment, a {100} texture is formed in a surface layer portion of an Fe-based metal sheet by using γ-α transformation accompanying decarburization (and further demanganization), and thereafter a ferrite-forming element is made to diffuse into a partial or whole decarburized region and further over the region from its surface, and at the time of cooling, the whole Fe-based metal sheet is made to be oriented in {100}.
This embodiment as above is based on the fact found by the present inventors that {100} crystal grains in the texture formed in the surface preferentially grow at an A3 point or higher in a heating process to be performed for the diffusion of the ferrite-forming element and further when the ferrite-forming element is made to diffuse into the inner portion to make the Fe-based metal sheet alloyed therewith and then cooling is performed, an accumulation degree of {200} planes in a sheet plane of the Fe-based metal sheet increases.
[Explanation of the Basic Principle of the Second Embodiment of the Present Invention]
First, the basic principle of this embodiment capable of obtaining a high accumulation degree of {200} planes will be explained based on FIG. 3A to FIG. 3D, by taking the case of decarburization as an example.
(a) Seeding of a Texture
When being decarburized until C becomes less than 0.02 mass %, the Fe-based metal sheet containing C: 0.02 mass % or more and having a composition of the α-γ transforming component is heated to a temperature at which a structure is turned into an α single phase and to a temperature of a γ single phase or a two-phase region of a γ phase and an α phase (namely, a temperature of an A1 point or higher) to decarburize the surface layer portion of the Fe-based metal sheet until C becomes less than 0.02 mass %. Thereby, the γ-α transformation is made to occur in a process of the decarburization to turn the decarburized surface layer portion into the α phase.
At this time, the decarburization progresses the most in the <110> direction in the γ phase having large gaps between lattices, and in this portion, the C concentration becomes less than 0.02% and the transformation to the α phase occurs. A {110} plane in the γ phase becomes the {100} plane when the γphase is turned into the α phase in a BCC structure, and thus in the α phase after the decarburization, the {100} plane is preferentially formed. Further, the growth, of the crystal grains in the α phase formed in the surface, in the sheet thickness direction is slow because its rate is controlled by a decarburization rate, and thus the crystal grains in the α phase formed in the surface grow in a direction parallel to the sheet plane. Further, in the surface of the metal sheet, the {100} plane preferentially grows by taking surface energy as driving force. As a result, the whole surface of the metal sheet becomes a structure oriented in {100} finally. By this process, as shown in FIG. 3A, a base metal sheet 1 having an inner region 4 made of Fe in the α phase and having the accumulation degree of the {200} planes in the decarburized region brought to 20% or more can be obtained. Further, a seed of crystal that satisfies the condition of the Z value is formed in the structure formed at the time of the decarburization by taking the surface energy as driving force.
(b) (Formation of a Second Layer)
Next, as shown in FIG. 3B, the ferrite-forming element such as Al is bonded to one surface or both surfaces of the base metal sheet 1 after the decarburization by using a vapor deposition method or the like to form a second layer 5.
(c) Saving of the Texture
Next, the base metal sheet 1 having had the ferrite-forming element bonded thereto is heated to the A3 point of the base metal sheet 1 to make the ferrite-forming element diffuse into the partial or whole decarburized region in the base metal sheet 1 to make the base metal sheet 1 alloyed therewith. Thereby, as shown in FIG. 3C, the α phase is formed in an alloyed region 6. Alternately, the ferrite-forming element is made to diffuse into the inner portion over the decarburized region to make the base metal sheet 1 alloyed therewith, and the alloyed region is turned to have the α single phase component partially, and thereby the region is turned into the α phase. At this time, the region is transformed while taking over orientation of the region formed by the decarburization, so that the structure oriented in {100} is formed also in the alloyed region 6. Further, the orientation in {100} is further increased even in crystal grains turned into the α phase previously. Further, when the ferrite-forming element is made to diffuse and the crystal is oriented, the seed of the crystal satisfying the condition of the Z value preferentially grows.
(d) Achievement of High Accumulation of the Texture
Next, the partially alloyed base metal sheet 1 is further heated to a temperature of not lower than the A3 point nor higher than 1300° C. and the temperature is held. The region of the α single phase component is an α-Fe phase not undergoing γ transformation, and thus the {100} crystal grains are maintained as they are, the {100} crystal grains preferentially grow in the region, and the accumulation degree of the {200} planes increases. Further, as shown in FIG. 3D, the region not having the α single phase component is transformed to the γ phase from the α phase.
Further, when a holding time of the temperature after the heating is prolonged, the {100} crystal grains are united to preferentially grow to large {100} crystal grains 7. As a result, the accumulation degree of the {200} planes further increases. Further, with the diffusion of Al, the region alloyed with Al is transformed to the α phase from the γ phase. At that time, in the region adjacent to the region to be transformed, crystal grains in the α phase oriented in {100} are already formed, and at the time of the transformation to the α phase from the γ phase, the region is transformed while taking over a crystal orientation of the adjacent crystal grains in the α phase. Thereby, the holding time is prolonged and the accumulation degree of the {200} planes increases.
(e) Growth of the texture
Next, the base metal sheet is cooled to a temperature of lower than the A 3point. At this time, as shown in FIG. 3E, a γ-Fe phase in an unalloyed inner region 10 is transformed to the α-Fe phase. This inner region 10 is adjacent to the region in which the crystal grains in the α phase oriented in {100} are already formed in a temperature region of the A3point or higher, and at the time of the transformation to the α phase from the γ phase, the inner region 10 is transformed while taking over the crystal orientation of the adjacent crystal grains in the α phase and larger crystal grains 9 in the α phase oriented in {100} are formed. Therefore, the accumulation degree of the {200} planes increases also in the region (see the state shown in FIG. 3E). By this phenomenon, the high accumulation degree of the {200} planes can be obtained even in the unalloyed region 10.
When at the stage of the preceding state shown in FIG. 3D, the temperature of the A3 point or higher is held until the whole metal sheet is alloyed, the structure having the high accumulation degree of the {200} planes is already formed in the whole metal sheet, and thus the cooling is performed while the state when the cooling is started is maintained.
Further, in the above explained example, the Fe-based metal sheet containing C: 0.02 mass % or more is used, but when an Fe-based metal sheet containing C: less than 0.02 mass % is used, carburization is performed before the decarburization to bring the C content in the region to be decarburized to 0.02 mass % or more.
In the above, the basic constitution of this embodiment was explained, and there will be further explained a limiting reason of each condition that defines a manufacturing method of this embodiment and preferable conditions of this embodiment.
[Fe-Based Metal to be the Base Material] (C Content)
In this embodiment, first, crystal grains oriented in {100} to serve as seeds for increasing the accumulation degree of the {200} planes are formed in the surface of the base metal sheet made of the Fe-based metal. Then, the γ-α transformation is made to progress in the metal sheet while taking over a crystal orientation of the crystal grains in the α phase to serve as the seeds finally, to thereby increase the accumulation degree of the {200} planes of the whole metal sheet.
In this embodiment, the seeds of the crystal grains oriented in {100} are formed in the surface of the base metal sheet by structure control using the γ-α transformation accompanying decarburization or demanganization. The Fe-based metal used for the base metal sheet has a composition of the α-γtransforming component, and the C content in the region to be decarburized is brought to 0.02 mass % or more.
Further, the Fe-based metal used for the base metal sheet has the α-γ transforming component, and the ferrite-forming element is made to diffuse into the metal sheet to make the metal sheet alloyed therewith, thereby making it possible to form a region having the α single phase based component. Further, the C content in the region to be decarburized is brought to 0.02 mass % or more, thereby making it possible to use the γ-αtransformation accompanying the decarburization.
For bringing the C content in the base metal sheet to 0.02 mass % or more, there is a method of using a base metal sheet manufactured from a molten material adjusted to contain C: 0.02 mass % or more by undergoing casting and rolling processes (a melting method). As another method, there is a method in which a base metal sheet having the C content of less than 0.02 mass % is used and in a surface layer portion of the base metal sheet, a region containing C: 0.02 mass % or more is formed by carburization.
In the case of the melting method, the range of the C content is set to not less than 0.02 mass % nor more than 1.0 mass %. When the C content is less than 0.02 mass %, it is not possible to use the formation of a {200} texture using the γ-α transformation accompanying the decarburization. Further, when the C content is more than 1.0% mass, a long time is required for the decarburization. The preferred range of the C content is not less than 0.05 mass % nor more than 0.5 mass %.
In the case of the carburization method, the range of the C content of the Fe-based metal of which the base metal sheet is made is set to 1 ppm or more and less than 0.02 mass %. Then, the surface layer of this Fe-based metal is subjected to the carburization so that the C concentration may become not less than 0.02 mass % nor more than 1.0 mass % in the same manner as that in the melting case.
Further, a carburizing range is set to a region down to a distance y from the surface, where the distance in a depth direction from the surface is set to y. This distance y is not less than 5 μm nor more than 50 μm. When the distance y is less than 5 μm, it is difficult to bring the accumulation degree of the {200} planes to 30% or more in the diffusion treatment after the decarburization, so that the distance y is set to 5 μm or more. Further, when the distance becomes greater than 50 μm, a long time is required for the carburization, and further a long time is required also for the decarburization of the whole carburized region. Further, an obtainable effect is also saturated, so that the preferred distance y is set to 50 μm or less. The carburizing method is not limited in particular, and a well-known gas carburizing method or the like may be performed.
Incidentally, the C content is preferably 0.005 mass % or less in terms of a magnetic property of a product metal sheet, so that in order to manufacture a steel sheet excellent in a magnetic property, silicon steel having the C content of 0.005 mass % or less is used to be subjected to carburization in a manner to have the above-described C concentration, which is advantageous for cost.
(Mn Content)
When Mn being an austenite stabilizing element is contained in the Fe-based metal, it is possible to form seeds of crystal grains oriented in {100} by structure control using the γ-α transformation accompanying demanganization. The demanganization is performed together with the decarburization, and thereby the surface layer portion is turned into the α phase more efficiently and the accumulation degree of the {200} planes in a decarburized and demanganized region is more increased. In order to exhibit such a function, the Mn content before performing the demanganization treatment is preferably set to 0.2 mass % or more.
The above-described structure control using the γ-α transformation can be performed even by the decarburization alone, so that Mn does not have to be contained. However, when Mn is contained, an effect of increasing electric resistance to decrease a core loss is also obtained, and thus Mn in a range of 2.0 mass % or less may also be contained according to need even when no demanganization is performed. From the above point, the range of the Mn content when Mn is contained is preferably set to 0.2 mass % to 2.0 mass %.
(Other Containing Elements)
In principle, being applicable to the Fe-based metal having the α-γ transforming component, this embodiment is not limited to the Fe-based metal in a specific composition range. Typical examples of the α-γ transforming component are pure iron, steel such as ordinary steel, and the like. For example, it is a component containing pure iron or steel containing C of 1 ppm to 0.10 mass % as described above or further containing Mn of 0.2 mass % to 2.0 mass % and a balance being composed of Fe and inevitable impurities as its base and containing an additive element as required. Instead, it may be silicon steel of the α-γ transforming component having C: 1.0 mass % or less and Si: 0.1 mass % to 2.5 mass % as its basic component. Further, as other impurities, a trace amount of Ni, Cr, Al, Mo, W, V, Ti, Nb, B, Cu, Co, Zr, Y, Hf, La, Ce, N, O, P, S, and/or the like are/is contained. Incidentally, Al and Mn are added to increase electric resistance, to thereby decrease a core loss and Co is added to increase the saturation magnetic flux density Bs, to thereby increase a magnetic flux density, which are also included in the present invention range.
(Thickness of the Base Metal Sheet)
The thickness of the base metal sheet is set to not less than 10 μm nor more than 6 mm. When the thickness is less than 10 μm, when the base metal sheets are stacked to be used as a magnetic core, the number of the sheets to be staked is increased to increase gaps, resulting in that a high magnetic flux density cannot be obtained. Further, when the thickness is greater than 6 mm, it is not possible to make the {100} texture grow sufficiently after cooling after the diffusion treatment, resulting in that a high magnetic flux density cannot be obtained.
[Decarburization Treatment]
In the decarburization treatment for turning the surface layer portion of the base metal sheet into the α phase, the base metal sheet is desirably heated in a decarburizing atmosphere to be decarburized in the following manner.
(Temperature of the Decarburization Treatment)
The temperature of the decarburization treatment is set to a temperature of the A1 point or higher and a temperature at which a structure is turned into an α single phase when the decarburization is performed until C becomes less than 0.02 mass %. The base metal sheet containing C: 0.02 mass % or more is heated to a temperature of a γ single phase or a two-phase region of a γ phase and an α phase (namely a temperature of the A1 point or higher) in order to make the γ-αtransformation occur by the decarburization.
(Atmosphere of the Decarburization Treatment)
With regard to the decarburizing atmosphere, a conventionally known method in manufacture of a grain-oriented electrical steel sheet can be employed. For example, there is a method in which decarburization is first performed in a weak decarburizing atmosphere, in a vacuum of 1 Torr or less, for example, or in a gas atmosphere of one type or two or more types of H2, He, Ne, Nr, Kr, Xe, Rn, and N2 at a temperature of lower than (a dew point—20)° C., and next decarburization is performed in a strong decarburizing atmosphere, or in a gas atmosphere in which an inert gas, or CO and CO2 is/are added to H2 at a temperature of (a dew point—20)° C. or higher, for example. In this case, if the decarburization is continued to the end in the weak decarburizing atmosphere, a long time is required.
(Period of Performing the Decarburization Treatment)
The period of performing the decarburization treatment is preferably not shorter than 0.1 minute nor longer than 600 minutes. When the period is shorter than 0.1 minute, it is difficult to bring the accumulation degree of the {200} planes to 20% or more after the decarburization, and when the period is long so as to exceed 600 minutes, too much cost is needed.
(Range of Performing the Decarburization Treatment)
The range of performing the decarburization treatment is a range down to a distance x, where the distance in the depth direction from the surface is set to x, and the distance x is not less than 5 μm nor more than 50 μm. When the distance x is less than 5 μm, it is difficult to bring the accumulation degree of the {200} planes to 30% or more in the diffusion treatment after the decarburization. For this reason, the distance x in the depth direction from the surface is set to 5 μm or more. Further, when the distance is greater than 50 μm, a long time is required for the decarburization, and further the accumulation degree of the {200} planes is saturated, and thus it is not advantageous industrially. Thus, the distance x is set to 50 μm or less.
(Other Decarburizing Methods)
Further, as described in Patent Literature 6, it is also possible that a material promoting decarburization is applied to a surface of a steel sheet as an annealing separating agent and this is wound around a coil and is subjected to coil annealing, to thereby form a decarburized region. Further, it is also possible that the above-described annealing separating agent is applied to a surface of a steel sheet in α single sheet form and the steel sheets are stacked to be subjected to annealing at the above-described temperature for a similar time, to thereby form a decarburized region.
(C Content after the Decarburization)
The C content after the decarburization is set to less than 0.02 mass % in order to obtain an α-phase single phase structure as described above. It is preferably 0.005 mass % or less in terms of the magnetic property of a product.
(Accumulation Degree of the {200} Planes after the Decarburization)
It is preferred that the accumulation degree of the {200} planes in the decarburized region after the decarburization should become 20% or more by performing the decarburization annealing under the above conditions. When the accumulation degree of the {200} planes is less than 20%, it is difficult to bring the accumulation degree of the {200} planes to 30% or more in the diffusion treatment to be performed subsequently. Further, the upper limit of the accumulation degree of the {200} planes is preferably set to 99%. When it is greater than 99%, the magnetic property deteriorates. The accumulation degree of the {200} planes is adjusted to fall within the above-described range by selecting the conditions of the decarburizing temperature, the decarburizing time, the decarburizing atmosphere, and the like. Incidentally, the measurement of the accumulation degree of the plane in the above-described orientation plane can be performed by X-ray diffraction using a MoKα ray similarly to the first embodiment.
[Demanganization Treatment]
In this embodiment, the decarburization treatment and the demanganization treatment may also be used in combination by containing Mn in the base metal sheet. The demanganization treatment is performed simultaneously with the decarburization or subsequently to the decarburization under the following conditions. Incidentally, as described in Patent Literature 6, it is also possible to perform the decarburization treatment and the demanganization treatment simultaneously in a state where steel sheets each have an annealing separating agent containing a material promoting decarburization and a material promoting demanganization applied thereto to be staked.
(Temperature and Range of the Demanganization Treatment)
The temperature at which the demanganization treatment is performed is set to a temperature of the A1 point or higher similarly to the decarburization. With regard to a demanganizing atmosphere, the demanganization treatment may be performed under a reduced pressure atmosphere. Further, the period of performing the demanganization treatment is preferably set to fall within a range of not shorter than 0.1 minute nor longer than 600 minutes similarly to the decarburization.
(Range of Performing the Demanganization Treatment)
The range of performing the demanganization treatment is a range down to a distance x, where the distance in the depth direction from the surface is set to x, and the distance x is preferably not less than 5 μm nor more than 50 μm. When the distance x is less than 5 μm, it is difficult to bring the accumulation degree of the {200} planes to 30% or more in the diffusion treatment after the demanganization. For this reason, the preferred distance x in the depth direction from the surface is set to 5 μm or more. Further, when the distance is greater than 50 μm, a long time is required for the demanganization, and further the accumulation degree of the {200} planes is saturated, and thus it is not advantageous industrially. Thus, the preferred distance x is set to 50 μm or less.
(Accumulation Degree of the {200} Planes after the Demanganization)
It is preferred that the accumulation degree of the {200} planes in the region having been subjected to the demanganization treatment should become 20% or more after the demanganization by performing the decarburization annealing under the above conditions. When the accumulation degree of the {200} planes is less than 20%, it is difficult to bring the accumulation degree of the {200} planes to 30% or more in the diffusion treatment to be performed subsequently. The upper limit of the accumulation degree of the {200} planes is preferably set to 99%. When it is greater than 99%, the magnetic property deteriorates.
[Different Metal]
Next, a different metal except Fe is made to diffuse into the base metal sheet having had the surface layer portion turned into the α phase by the decarburization to increase the region of the {100} texture in the thickness direction of the metal sheet. As the different metal, the ferrite-forming element is used. As a procedure, first, the different metal is bonded in a layered form as the second layer to one surface or both surfaces of the base metal sheet made of the Fe-based metal of the α-γ transforming component. Then, a region alloyed by having had elements of the different metal diffuse thereinto is turned to have the α single phase based component and to be able to be maintained as not only the region having been subjected to the decarburization (or further the demanganization) to be transformed to the α phase, but also a seed oriented in {100} for increasing the accumulation degree of the {200} planes in the metal sheet. As such a ferrite-forming element, at least one type of Al, Cr, Ga, Ge, Mo, Sb, Si, Sn, Ta, Ti, V, W, and Zn can be used alone or in a combined manner.
As a method of bonding the different metal in a layered form to the surface of the base metal sheet, there can be employed various methods such as a plating method of hot dipping, electrolytic plating, or the like, a rolling clad method, a dry process of PVD, CVD, or the like, and further powder coating. As a method of efficiently bonding the different metal for industrially implementing the method, the plating method or the rolling clad method is suitable.
The thickness of the different metal before the heating when the different metal is bonded is preferably not less than 0.05 μm nor more than 1000 μm. When the thickness is less than 0.05 μm, it is not possible to obtain the sufficient accumulation degree of the {200} planes. Further, when the thickness exceeds 1000 μm, even when the different metal layer is made to remain, its thickness becomes larger than necessary.
[Heating and Diffusion Treatment]
The base metal sheet having had the ferrite-forming element bonded thereto is heated up to the A3 point of the base metal sheet, to thereby make the ferrite-forming element diffuse into the partial or whole region in the base metal sheet to make the base metal sheet alloyed therewith. The α phase is maintained in the region alloyed with the ferrite-forming element. Alternately, the ferrite-forming element is made to diffuse into the inner portion over the decarburized region to make the base metal sheet alloyed therewith, and the alloyed region is turned to have the α single phase component partially, and thereby the region is turned into the α phase. At this time, the region is transformed while taking over the orientation of the region formed by the decarburization, so that the accumulation degree of the {200} planes further increases. As a result, in the alloyed region, a structure in which the accumulation degree of the {200} planes in the α-Fe phase becomes not less than 25% nor more than 50% and in accordance with it, the accumulation degree of the {200} planes in the α-Fe phase becomes not less than 1% nor more than 40% is formed.
Then, the base metal sheet is further heated to a temperature of not lower than the A3 point nor higher than 1300° C. and the temperature is held. The region alloyed already is turned into an α single phase structure that is not transformed to the γphase, so that the {100} crystal grains are maintained as they are, and in the region, the crystal grains in the {100} texture preferentially grow and the accumulation degree of the {200} planes increases. Further, the region not having the α single phase component is transformed to the γ phase.
Further, when the holding time is prolonged, the crystal grains in the {100} texture are united to one another to preferentially grow. As a result, the accumulation degree of the {200} planes further increases. Further, with the further diffusion of the ferrite-forming element, the region alloyed with the ferrite-forming element is transformed to the α phase from the γ phase. At this time, as shown in FIG. 4A, in the regions adjacent to the regions to be transformed, crystal grains 7 in the α phase oriented in {100} are already formed, and at the time of the transformation to the α phase from the γ phase, the regions alloyed with the ferrite-forming element are transformed while taking over a crystal orientation of the adjacent crystal grains 7 in the α phase. Thereby, the holding time is prolonged and the accumulation degree of the {200} planes increases. Further, as a result, the accumulation degree of the {200} planes decreases.
Incidentally, in order to finally obtain the high accumulation degree of the {200} planes of 50% or more, it is preferred that the holding time should be adjusted to, at this stage, bring the accumulation degree of the {200} planes in the α-Fe phase to 30% or more and bring the accumulation degree of the {200} planes in the α-Fe phase to 30% or less. Further, when the A3 point or higher is held until the whole metal sheet is alloyed, as shown in FIG. 4C, the α single phase structures are formed up to the center portion of the metal sheet and grain structures oriented in {100} reach the center of the metal sheet.
A holding temperature after the temperature is increased is set to not lower than A3 point nor higher than 1300° C. Even when the metal sheet is heated at a temperature higher than 1300° C., an effect with respect to the magnetic property is saturated. Further, cooling may be started immediately after the temperature reaches the holding temperature, or cooling may also be started after the temperature is held for 6000 minutes or shorter. When this condition is satisfied, the achievement of high accumulation of the seeds oriented in the {200} plane further progresses to make it possible to more securely bring the accumulation degree of the {200} planes in the α-Fe phase to 30% or more after the cooling.
[Cooling after the Heating and Diffusion Treatment]
After the diffusion treatment, when the cooling is performed while the region that is not alloyed with the ferrite-forming element is remaining, as shown in FIG. 4B, at the time of the transformation to the α phase from the γ phase, the unalloyed region is transformed while taking over the crystal orientation of the regions in which the crystal grains 9 in the α phase oriented in {100} are already formed. Thereby, the accumulation degree of the {200} planes increases, and the metal sheet having the texture in which the accumulation degree of the {200} planes in the α-Fe phase is not less than 30% nor more than 99% and the accumulation degree of the {200} planes in the α-Fe phase is not less than 0.01% nor more than 30% is obtained, the crystal satisfying the condition of the Z value grows, and a high magnetic flux density can be obtained in an arbitrary direction in the metal sheet plane.
Further, as shown in FIG. 4C, when the A3 point or higher is held until the whole metal sheet is alloyed and the grain structures oriented in {100} reach the center of the metal sheet, as shown in FIG. 4D, the metal sheet is cooled as it is, and the texture in which the grain structures oriented in {100} reach the center of the metal sheet can be obtained. Thereby, the whole metal sheet is alloyed with the different metal, and the metal sheet having the texture in which the accumulation degree of the {200} planes in the α-Fe phase is not less than 30% nor more than 99% and the accumulation degree of the {200} planes in the α-Fe phase is not less than 0.01% nor more than 30% is obtained.
As above, the value of the accumulation degree of the {200} planes and the remaining state of the different metal on the surface of the base metal sheet change depending on the holding time of the temperature of the A3 point or higher and the holding temperature. The example shown in FIG. 4B is in a state where the grain structures oriented in {100} do not reach up to the center of the metal sheet and the different metal also remains on the surfaces, but it is also possible to obtain the grain structures oriented in {100} up to the center of the metal sheet and to alloy all the second layers on the surfaces.
Incidentally, at the time of the cooling after the diffusion treatment, a cooling rate is preferably not less than 0.1° C./sec nor more than 500° C./sec. When the cooling rate is less than 0.1° C./sec, a long time is required for the cooling, which is not appropriate, and when the cooling rate is greater than 500° C./sec, the metal sheet is sometimes deformed, and thus the cooling rate is preferably 500° C./sec or less.
Incidentally, when the second layers are made to remain on the obtainable Fe-based metal sheet having a thickness of not less than 10 μm nor more than 6 mm, the thickness of the second layer is preferably set to not less than 0.01 μm nor more than 500 μm. Further, a ratio of the α single phase region alloyed at this stage is preferably 1% or more in a cross section of the Fe-based metal sheet.
Further, it is also possible to form a structure as shown in FIG. 5, and in this case, an average cooling rate is set to satisfy the condition similar to that of the first embodiment, and thereby the above can be achieved.
EXAMPLE
Next, there will be explained experiments conducted by the present inventors. Conditions and the like in these experiments are examples employed for confirming the applicability and effects of the present invention, and the present invention is not limited to these examples.
Example 1
In this example, base metal sheets of No. 1 to No. 17 each made of a component A or B shown in Table 1 below were manufactured under various rolling conditions, to then have various different metals applied thereto as a second layer, and then Fe-based metal sheets were fabricated, of which the previously described Z value (=(A +0.97B)/0.98C) and the magnetic flux density difference ΔB were examined. Further, the relationship between various manufacturing conditions and an accumulation degree of {200} planes was also examined. Further, effects obtained by changing a starting temperature in an a-region rolling process were also examined in detail.









TABLE 1







COMPONENT
A3
ELEMENT MASS %

















SERIES
POINT
C
Si
Mn
Al
P
N
S
0
OTHER




















A
925
0.0008
0.3
0.3
0.5
0.0003
0.0002
<0.0004
0.0002
—


B
1010
0.0012
1.1
0.8
0.1
0.0002
0.0003
<0.0004
0.0001
—


C
915
0.0032
0.2
0.08
0.05
0.0001
0.0003
<0.0004
0.0001
—


D
870
0.0041
0.1
1.5
0.2
0.0001
0.0002
<0.0004
0.0001
—


E
942
0.0105
0.2
0.5
0.7
0.0001
0.0003
<0.0004
0.0001
Cr: 0.5









First, ingots each having the component A or B shown in Table 1 and a balance being composed of Fe and inevitable impurities were melted by vacuum melting. Then, these were used as rolling materials to be worked into cold-rolled sheets (the base metal sheets) each having a predetermined thickness under conditions of hot rolling, α-region rolling, and cold rolling shown in Table 2 below.









TABLE 2









α-REGION










HOT ROLLING
ROLLING












START
FINISH
START
FINISH

















BASE


TEMPER-
THICK-
TEMPER-
THICK-
REDUC-
TEMPER-
THICK-
TEMPER-


MATERIAL

A3
ATURE
NESS
ATURE
NESS
TION
ATURE
NESS
ATURE


No.
COMPONENT
POINT
° C.
mm
° C.
mm
RATIO
° C.
mm
° C.





1
A
925
1150
250
1000
10
−3.22
950
10
920


2
A
925
1150
250
1000
10
−3.22
920
10
830


3
A
925
1150
250
1000
10
−3.22
850
10
830


4
A
925
1150
250
1000
10
−3.22
750
10
730


5
A
925
1150
250
1000
10
−3.22
650
10
640


6
A
925
1150
250
1000
10
−3.22
550
10
540


7
A
925
1150
250
1000
10
−3.22
450
10
450


8
A
925
1150
250
1000
10
−3.22
300
10
350


9
A
925
1150
250
1000
10
−3.22
250
10
250


10
B
1010
1200
280
1050
50
−1.72
1050
50
980


11
B
1010
1200
280
1050
50
−1.72
950
50
880


12
B
1010
1200
280
1050
50
−1.72
850
50
770


13
B
1010
1200
280
1050
50
−1.72
750
50
660


14
B
1010
1200
280
1050
50
−1.72
600
50
580


15
B
1010
1200
280
1050
50
−1.72
450
50
485


16
B
1010
1200
280
1050
50
−1.72
300
50
390


17
B
1010
1200
280
1050
50
−1.72
250
50
230













α-REGION












ROLLING
COLD ROLLING














FINISH
START
FINISH
TOTAL
REDUCTION
















BASE
THICK-
REDUC-
THICK-
THICK-
REDUC-
REDUC-
RATIO OF α



MATERIAL
NESS
TION
NESS
NESS
TION
TION
REGION +



No.
min
RATIO
mm
mm
RATIO
RATIO
COLD ROLLING







1
2.5
−1.39
2.5
0.2
−2.53
−7.13
−3.91



2
2.5
−1.39
2.5
0.2
−2.53
−7.13
−3.91



3
2.5
−1.39
2.5
0.2
−2.53
−7.13
−3.91



4
2.5
−1.39
2.5
0.2
−2.53
−7.13
−3.91



5
2.5
−1.39
2.5
0.2
−2.53
−7.13
−3.91



6
2.5
−1.39
2.5
0.2
−2.53
−7.13
−3.91



7
2.5
−1.39
2.5
0.2
−2.53
−7.13
−3.91



8
2.5
−1.39
2.5
0.2
−2.53
−7.13
−3.91



9
2.5
−1.39
2.5
0.2
−2.53
−7.13
−3.91



10
3.0
−2.81
3.0
0.5
−1.79
−6.33
−4.61



11
3.0
−2.81
3.0
0.5
−1.79
−6.33
−4.61



12
3.0
−2.81
3.0
0.5
−1.79
−6.33
−4.61



13
3.0
−2.81
3.0
0.5
−1.79
−6.33
−4.61



14
3.0
−2.81
3.0
0.5
−1.79
−6.33
−4.61



15
3.0
−2.81
3.0
0.5
−1.79
−6.33
−4.61



16
3.0
−2.81
3.0
0.5
−1.79
−6.33
−4.61



17
3.0
−2.81
3.0
0.5
−1.79
−6.33
−4.61










In the case of the component A, the ingots each having a thickness of 250 mm heated to 1150° C. were first subjected to hot rolling at a reduction ratio of −3.22 in terms of true strain, and hot-rolled sheets each having a thickness of 10 mm were obtained. Next, these hot-rolled sheets were each subjected to α-region rolling at a reduction ratio of −1.39 in terms of true strain at a temperature of 300 to 1000° C. These rolled sheets obtained by the α-region rolling were pickled, and then the base metal sheets were obtained by cold rolling. At this time, the reduction ratio was −2.53 in terms of true strain, and as a result, the thickness of each of the obtained base metal sheets was 0.2 mm.
In the case of the component B, the ingots each having a thickness of 280 mm heated to 1200° C. were first subjected to hot rolling at a reduction ratio of −1.72 in terms of true strain, and hot-rolled sheets each having a thickness of 50 mm were obtained. Next, these hot-rolled sheets were each subjected to α-region rolling at a reduction ratio of −2.81 in terms of true strain at a temperature of 300 to 1050° C. These rolled sheets obtained by the α-region rolling were pickled, and then the base metal sheets were obtained by cold rolling. At this time, the reduction ratio was −1.79 in terms of true strain, and as a result, the thickness of each of the obtained base metal sheets was 0.5 mm.
With respect to the base metal sheets obtained by the above procedure, a texture of a surface layer portion of each of the base materials was measured by X-ray diffraction to obtain an accumulation degree of {200} planes and an accumulation degree of {222} planes by the previously described method. Further, thinning was performed so that a structure could be observed from a direction perpendicular to an L cross section, and a region up to 1⁄4t (t represents a thickness) from the surface was observed. The main phase of each of the obtained base metal sheets at room temperature was an α-Fe phase. Further, as a result of measurement, the A3 point at which the α-γ transformation occurred was 925° C. in the component A and 1010° C. in the component B.
Next, both surfaces of each of the base metal sheets of No. 1 to No. 17 shown in Table 2 were coated with each of various different metal elements as the second layer by a vapor deposition method, a sputtering method, or an electroplating method. As shown in Table 3 and Table 4 below, as the different metal element, any one of Al, Si, Mo, Ga, Sn, Ti, Ge, Sb, V, and W was selected. The thickness of each of the coatings was as shown in Table 3 and Table 4.
Next, an experiment was performed in which a heat treatment was performed on the base metal sheets to each of which the second layers were bonded under various conditions. A gold image furnace was used for the heat treatment, and a holding time was controlled by program control. During which the temperature increased to be held, the heat treatment was performed in an atmosphere vacuumed to a pressure of 10−3 Pa level. At the time of cooling, in the case of a cooling rate of 1° C./sec or lower, temperature control was performed in a vacuum by furnace output control. Further, in the case of the cooling rate of 10° C./sec or more, an Ar gas was introduced and the cooling rate was controlled by adjustment of its flow rate.
Here, there was examined a change in the texture among a temperature increasing process of heating up to the A3 point, a holding process of heating to a temperature of not lower than the A3 point nor higher than 1300° C. and holding the temperature, and a cooling process of cooling to a temperature of lower than the A3 point. Specifically, three base metal sheets with the same combination of the base material-coating conditions were prepared, of which a change in the texture was examined by performing a heat treatment experiment in each of the processes.
A sample for the temperature increasing process was fabricated in such a manner that the base metal sheet was heated from room temperature to the A3 point at a predetermined temperature increasing rate and was cooled to room temperature without any holding time. The cooling rate was set to 100° C./sec. The texture was measured by the method using the previously described X-ray diffraction method, and the X-ray was emitted from its surface, and the accumulation degree of {200} planes in the α-Fe phase and the accumulation degree of {222} planes in the α-Fe phase were obtained in an inverse pole figure.
A sample for the holding process was fabricated in such a manner that the base metal sheet was heated from room temperature to a predetermined temperature over the A3 point at a predetermined temperature increasing rate and was cooled to room temperature after a predetermined holding time. Then, the texture of the fabricated sample was measured in the same manner, and the accumulation degrees of {200} and {222} planes in the α-Fe phase were obtained.
A sample for the cooling process was fabricated in such a manner that the base metal sheet was heated from room temperature to a predetermined temperature over the A3 point at a predetermined temperature increasing rate and was cooled to room temperature at a predetermined cooling rate after a predetermined holding time. Further, in order to evaluate the accumulation degrees of {200} and {222} planes at an unalloyed position, a test piece was fabricated by removing a layer from the surface of the fabricated sample to a predetermined distance so that the unalloyed position might become an evaluation surface. Incidentally, when the whole metal sheet was alloyed, the evaluation surface was set to a position of 1⁄2 of the sheet thickness. With regard to the measurement of the texture of the fabricated sample, the X-ray was emitted from the surface of the test piece and from a predetermined surface of the test piece from which the layer was removed, and the accumulation degrees of {200} and {200} planes in the α-Fe phase of the surfaces were obtained in the same manner.
Next, magnetometry was performed in order to evaluate obtained products. First, the average magnetic flux density B50 to a magnetizing force of 5000 A/m and the magnetic flux density difference ΔB were obtained by using a SST (Single Sheet Tester). At this time, a measurement frequency was set to 50 Hz. When the average magnetic flux density B50 was obtained, as shown in FIG. 1, the magnetic flux density B50 was obtained every 22.5° in a circumferential direction of the product and an average value of the magnetic flux densities B50 in these 16 directions was calculated. Further, of the magnetic flux densities B50 in these 16 directions, the difference between the maximum value and the minimum value was set to the magnetic flux density difference ΔB. Next, the saturation magnetic flux density Bs was obtained by using a VSM (Vibrating Sample Magnetometer). The magnetizing force applied at this time was 0.8×106 A/m. An evaluation value was set to the ratio B50/Bs of the average magnetic flux density B50 to the saturation magnetic flux density.
Further, by the previously described X-ray diffraction, intensity ratios of {001}<470>, {116}<6 12 1>, and {223}<692> were calculated, and thereby the previously described Z value was calculated.
Table 3 and Table 4 below show the accumulation degrees of the {200} planes and the accumulation degrees of the {222} planes measured in the respective processes during the manufacture and after the manufacture, the Z values of the obtained Fe-based metal sheets, and evaluation results of the magnetometry.









TABLE 3









MANUFACTURE









SEEDING












BASE

SEEDING
SEEDING














MATERIAL
SECOND LAYER
TEMPERATURE

ACCUMULATION
ACCUMULATION


















THICK-

THICK-

INCREASING

DEGREE OF
DEGREE OF




NESS

NESS
BONDING
RATE
MEASURED
{200} PLANES
{222} PLANES


No.
No.
mm
ELEMENT
μm
METHOD
° C./sec
TEMPERATURE
IN αFe
IN αFe





1
1
0.2
Al
4
VAPOR
20
925
14
13







DEPOSITION


2
2
0.2
Al
4
VAPOR
20
925
19
12







DEPOSITION


3
3
0.2
Al
4
VAPOR
20
925
25
10







DEPOSITION


4
4
0.2
Al
4
VAPOR
20
925
27
9.6







DEPOSITION


5
5
0.2
Al
4
VAPOR
20
925
33
9.1







DEPOSITION


6
6
0.2
Al
4
VAPOR
20
925
34
8.8







DEPOSITION


7
7
0.2
Al
4
VAPOR
20
925
35
8.7







DEPOSITION


8
8
0.2
Al
4
VAPOR
20
925
28
9.5







DEPOSITION


9
9
0.2
Al
4
VAPOR
20
925
27
9.8







DEPOSITION


10
1
0.2
Si
5
SPUTTERING
70
925
14
13


11
2
0.2
Si
5
SPUTTERING
70
925
18
12


12
3
0.2
Si
5
SPUTTERING
70
925
26
10


13
4
0.2
Si
5
SPUTTERING
70
925
27
9.4


14
5
0.2
Si
5
SPUTTERING
70
925
32
8.8


15
6
0.2
Si
5
SPUTTERING
70
925
35
8.2


16
7
0.2
Si
5
SPUTTERING
70
925
35
8.3


17
8
0.2
Si
5
SPUTTERING
70
925
29
9.7


18
9
0.2
Si
5
SPUTTERING
70
925
28
10


19
1
0.2
Mo
1.5
SPUTTERING
10
925
15
12


20
2
0.2
Mo
1.5
SPUTTERING
10
925
17
11


21
3
0.2
Mo
1.5
SPUTTERING
10
925
26
9.8


22
4
0.2
Mo
1.5
SPUTTERING
10
925
28
9.2


23
5
0.2
Mo
1.5
SPUTTERING
10
925
33
8.6


24
6
0.2
Mo
1.5
SPUTTERING
10
925
35
8.3


25
7
0.2
Mo
1.5
SPUTTERING
10
925
36
8.1


26
8
0.2
Mo
1.5
SPUTTERING
10
925
27
10


27
9
0.2
Mo
1.5
SPUTTERING
10
925
26
11


28
1
0.2
Ga
3
VAPOR
0.5
925
15
13







DEPOSITION


29
2
0.2
Ga
3
VAPOR
0.5
925
17
12







DEPOSITION


30
3
0.2
Ga
3
VAPOR
0.5
925
26
10







DEPOSITION


31
4
0.2
Ga
3
VAPOR
0.5
925
28
9.3







DEPOSITION


32
5
0.2
Ga
3
VAPOR
0.5
925
34
8.7







DEPOSITION


33
6
0.2
Ga
3
VAPOR
0.5
925
35
8.1







DEPOSITION


34
7
0.2
Ga
3
VAPOR
0.5
925
35
7.8







DEPOSITION


35
8
0.2
Ga
3
VAPOR
0.5
925
27
10







DEPOSITION


36
9
0.2
Ga
3
VAPOR
0.5
925
25
11







DEPOSITION


37
1
0.2
Sn
6
ELECTROLYTiC
5
925
16
13







PLATING


38
2
0.2
Sn
6
ELECTROLYTIC
5
925
19
11







PLATING


39
3
0.2
Sn
6
ELECTROLYTIC
5
925
27
9.5







PLATING


40
4
0.2
Sn
6
ELECTROLYTIC
5
925
28
9.1







PLATING


41
5
0.2
Sn
6
ELECTROLYTIC
5
925
32
8.7







PLATING


42
6
0.2
Sn
6
ELECTROLYTIC
5
925
33
8.3







PLATING


43
7
0.2
Sn
6
ELECTROLYTIC
5
925
34
8.2







PLATING


44
8
0.2
Sn
6
ELECTROLYTIC
5
925
27
9.7







PLATING


45
9
0.2
Sn
6
ELECTROLYTIC
5
925
26
11







PLATING












MANUFACTURE










SAVING AND ACHIEVEMENT OF




HIGH ACCUMULATION
















ACHIEVEMENT
ACHIEVEMENT






OF HIGH
OF HIGH











ACCUMULATION
ACCUMULATION
GROWTH













ACCUMULATION
ACCUMULATION

1⁄2t
1⁄2t















HOLDING
HOLDING
DEGREE OF
DEGREE OF
COOLING
ACCUMULATION
ACCUMULATION



TEMPERATURE
TIME
{200} PLANES
{222} PLANES
RATE
DEGREE OF
DEGREE OF


No.
° C.
sec
IN αFe
IN αFe
° C./sec
{200} PLANES
{222} PLANES





1
1000
20
16
13
150
16
13


2
1000
20
25
10.4
150
25
10.4


3
1000
20
30
9.1
150
30
9.1


4
1000
20
41
3.4
150
41
3.4


5
1000
20
53
1.8
150
53
1.8


6
1000
20
52
2.1
150
52
2.1


7
1000
20
50
2.3
150
50
2.3


8
1000
20
38
3.8
150
38
3.8


9
1000
20
37
4.2
150
37
4.2


10
1050
10
17
12
250
17
12


11
1050
10
24
12
250
24
12


12
1050
10
31
8.7
250
31
8.7


13
1050
10
42
3.2
250
42
3.2


14
1050
10
55
1.7
250
55
1.7


15
1050
10
54
1.9
250
54
1.9


16
1050
10
51
2.2
250
51
2.2


17
1050
10
39
4.1
250
39
4.1


18
1050
10
37
4.5
250
37
4.5


19
1250
10
15
13
10
15
13


20
1250
10
23
13
10
23
13


21
1250
10
30
9.3
10
30
9.3


22
1250
10
41
4.1
10
41
4.1


23
1250
10
52
2.4
10
52
2.4


24
1250
10
52
2.6
10
52
2.6


25
1250
10
51
2.9
10
51
2.9


26
1250
10
38
4.8
10
38
4.8


27
1250
10
37
5.5
10
37
5.5


28
980
100
17
11
50
17
11


29
980
100
27
9.8
50
27
9.8


30
980
100
33
8.5
50
33
8.5


31
980
100
43
3.5
50
43
3.5


32
980
100
57
1.8
50
57
1.8


33
980
100
56
2.1
50
56
2.1


34
980
100
55
2.3
50
55
2.3


35
980
100
40
3.9
50
40
3.9


36
980
100
37
4.5
50
37
4.5


37
1100
20
16
13
350
16
13


38
1100
20
27
11
350
27
11


39
1100
20
32
9.4
350
32
9.4


40
1100
20
45
3.1
350
45
3.1


41
1100
20
58
1.4
350
58
1.4


42
1100
20
57
1.9
350
57
1.9


43
1100
20
56
2.1
350
56
2.1


44
1100
20
43
3.8
350
43
3.8


45
1100
20
41
5.1
350
41
5.1













PRODUCT











TEXTURE EVALUATION












ACCUMULATION
ACCUMULATION

















DEGREE OF
DEGREE OF

MAGNETIC






{200} PLANES
{222} PLANES

FLUX DENSITY

α-REGION


















No.
IN αFe
IN αFe
Z
B50 T
Bs T
B50/Bs
ΔB T
NOTE
TEMPERATURE







1
16
13
1.2
1.60
2.05
0.78
0.070
COMPARATIVE
950











EXAMPLE 1



2
25
10.4
2.1
1.66
2.05
0.81
0.065
PRESENT
920











INVENTION











EXAMPLE 1



3
30
9.1
5.8
1.71
2.05
0.83
0.060
PRESENT
850











INVENTION











EXAMPLE 2



4
41
3.4
23
1.77
2.05
0.86
0.056
PRESENT
750











INVENTION











EXAMPLE 3



5
53
1.8
160
1.84
2.05
0.90
0.018
PRESENT
650











INVENTION











EXAMPLE 4



6
52
2.1
120
1.87
2.05
0.91
0.021
PRESENT
550











INVENTION











EXAMPLE 5



7
50
2.3
42
1.86
2.05
0.91
0.070
PRESENT
450











INVENTION











EXAMPLE 6



8
38
3.8
3.5
1.80
2.05
0.88
0.145
PRESENT
300











INVENTION











EXAMPLE 7



9
37
4.2
1.1
1.78
2.05
0.87
0.220
COMPARATIVE
250











EXAMPLE 2



10
17
12
1.4
1.60
2.05
0.78
0.080
COMPARATIVE
950











EXAMPLE 3



11
24
12
2.5
1.65
2.05
0.80
0.074
PRESENT
920











INVENTION











EXAMPLE 8



12
31
8.7
3.8
1.66
2.05
0.81
0.070
PRESENT
850











INVENTION











EXAMPLE 9



13
42
3.2
27
1.79
2.05
0.87
0.054
PRESENT
750











INVENTION











EXAMPLE 10



14
55
1.7
156
1.88
2.05
0.92
0.015
PRESENT
650











INVENTION











EXAMPLE 11



15
54
1.9
134
1.87
2.05
0.91
0.025
PRESENT
550











INVENTION











EXAMPLE 12



16
51
2.2
51
1.86
2.05
0.91
0.036
PRESENT
450











INVENTION











EXAMPLE 13



17
39
4.1
5.8
1.81
2.05
0.88
0.145
PRESENT
300











INVENTION











EXAMPLE 14



18
37
4.5
1.7
1.74
2.05
0.85
0.210
COMPARATIVE
250











EXAMPLE 4



19
15
13
1.3
1.59
2.05
0.78
0.087
COMPARATIVE
950











EXAMPLE 5



20
23
13
2.4
1.66
2.05
0.81
0.081
PRESENT
920











INVENTION











EXAMPLE 15



21
30
9.3
5.8
1.72
2.05
0.84
0.080
PRESENT
850











INVENTION











EXAMPLE 16



22
41
4.1
19
1.78
2.05
0.87
0.074
PRESENT
750











INVENTION











EXAMPLE 17



23
52
2.4
149
1.86
2.05
0.91
0.021
PRESENT
650











INVENTION











EXAMPLE 18



24
52
2.6
174
1.86
2.05
0.91
0.018
PRESENT
550











INVENTION











EXAMPLE 19



25
51
2.9
39
1.86
2.05
0.91
0.093
PRESENT
450











INVENTION











EXAMPLE 20



26
38
4.8
3.1
1.77
2.05
0.86
0.138
PRESENT
300











INVENTION











EXAMPLE 21



27
37
5.5
1.1
1.76
2.05
0.86
0.190
COMPARATIVE
250











EXAMPLE 6



28
17
11
1.2
1.61
2.05
0.79
0.082
COMPARATIVE
950











EXAMPLE 7



29
27
9.8
2.5
1.65
2.05
0.80
0.073
PRESENT
920











INVENTION











EXAMPLE 22



30
33
8.5
8.5
1.73
2.05
0.84
0.073
PRESENT
850











INVENTION











EXAMPLE 23



31
43
3.5
34
1.78
2.05
0.87
0.064
PRESENT
750











INVENTION











EXAMPLE 24



32
57
1.8
112
1.87
2.05
0.91
0.017
PRESENT
650











INVENTION











EXAMPLE 25



33
56
2.1
110
1.88
2.05
0.92
0.013
PRESENT
550











INVENTION











EXAMPLE 26



34
55
2.3
74
1.87
2.05
0.91
0.087
PRESENT
450











INVENTION











EXAMPLE 27



35
40
3.9
2.1
1.76
2.05
0.86
0.139
PRESENT
300











INVENTION











EXAMPLE 28



36
37
4.5
0.6
1.74
2.05
0.85
0.210
COMPARATIVE
250











EXAMPLE 8



37
16
13
0.8
1.60
2.05
0.78
0.086
COMPARATIVE
950











EXAMPLE 9



38
27
11
2.2
1.65
2.05
0.80
0.079
PRESENT
920











INVENTION











EXAMPLE 29



39
32
9.4
8.2
1.73
2.05
0.84
0.079
PRESENT
850











INVENTION











EXAMPLE 30



40
45
3.1
34
1.81
2.05
0.88
0.065
PRESENT
750











INVENTION











EXAMPLE 31



41
58
1.4
158
1.87
2.05
0.91
0.013
PRESENT
650











INVENTION











EXAMPLE 32



42
57
1.9
189
1.89
2.05
0.92
0.009
PRESENT
550











INVENTION











EXAMPLE 33



43
56
2.1
48
1.88
2.05
0.92
0.091
PRESENT
450











INVENTION











EXAMPLE 34



44
43
3.8
2.7
1.77
2.05
0.86
0.136
PRESENT
300











INVENTION











EXAMPLE 35



45
41
5.1
1.4
1.76
2.05
0.86
0.192
COMPARATIVE
250











EXAMPLE 10



















TABLE 4









MANUFACTURE









SEEDING












BASE

SEEDING
SEEDING














MATERIAL
SECOND LAYER
TEMPERATURE

ACCUMULATION
ACCUMULATION


















THICK-

THICK-

INCREASING

DEGREE OF
DEGREE OF




NESS

NESS
BONDING
RATE
MEASURED
{200} PLANES
{222} PLANES


No.
No.
mm
ELEMENT
μm
METHOD
° C./sec
TEMPERATURE
IN αFe
IN αFe





46
10
0.5
Ti
10
SPUTTERING
50
1010
14
13


47
11
0.5
Ti
10
SPUTTERING
50
1010
25
11


48
12
0.5
Ti
10
SPUTTERING
50
1010
30
9.3


49
13
0.5
Ti
10
SPUTTERING
50
1010
36
7.3


50
14
0.5
Ti
10
SPUTTERING
50
1010
38
6.7


51
15
0.5
Ti
10
SPUTTERING
50
1010
38
6.4


52
16
0.5
Ti
10
SPUTTERING
50
1010
27
9.5


53
17
0.5
Ti
10
SPUTTERING
50
1010
26
10


54
10
0.5
Ge
12
SPUTTERING
100
1010
13
14


55
11
0.5
Ge
12
SPUTTERING
100
1010
26
11


56
12
0.5
Ge
12
SPUTTERING
100
1010
30
9.1


57
13
0.5
Ge
12
SPUTTERING
100
1010
34
8.1


58
14
0.5
Ge
12
SPUTTERING
100
1010
37
6.9


59
15
0.5
Ge
12
SPUTTERING
100
1010
37
6.7


60
16
0.5
Ge
12
SPUTTERING
100
1010
26
10


61
17
0.5
Ge
12
SPUTTERING
100
1010
25
12


62
10
0.5
Sb
15
SPUTTERING
1
1010
12
15


63
11
0.5
Sb
15
SPUTTERING
1
1010
26
10


64
12
0.5
Sb
15
SPUTTERING
1
1010
30
9.1


65
13
0.5
Sb
15
SPUTTERING
1
1010
35
8


66
14
0.5
Sb
15
SPUTTERING
1
1010
37
7.2


67
15
0.5
Sb
15
SPUTTERING
1
1010
37
7.4


68
16
0.5
Sb
15
SPUTTERING
1
1010
26
10


69
17
0.5
Sb
15
SPUTTERING
1
1010
25
11


70
10
0.5
V
18
SPUTTERING
300
1010
13
13


71
11
0.5
V
18
SPUTTERING
300
1010
26
10


72
12
0.5
V
18
SPUTTERING
300
1010
31
8.5


73
13
0.5
V
18
SPUTTERING
300
1010
36
7.7


74
14
0.5
V
18
SPUTTERING
300
1010
38
6.3


75
15
0.5
V
18
SPUTTERING
300
1010
37
6.7


76
16
0.5
V
18
SPUTTERING
300
1010
27
10


77
17
0.5
V
18
SPUTTERING
300
1010
26
11


78
10
0.5
W
10
SPUTTERING
50
1010
13
14


79
11
0.5
W
10
SPUTTERING
50
1010
26
11


80
12
0.5
W
10
SPUTTERING
50
1010
31
9.1


81
13
0.5
W
10
SPUTTERING
50
1010
35
7.9


82
14
0.5
W
10
SPUTTERING
50
1010
37
6.4


83
15
0.5
W
10
SPUTTERING
50
1010
38
6.4


84
16
0.5
W
10
SPUTTERING
50
1010
26
10


85
17
0.5
W
10
SPUTTERING
50
1010
25
11












MANUFACTURE










SAVING AND ACHIEVEMENT OF




HIGH ACCUMULATION
















ACHIEVEMENT
ACHIEVEMENT






OF HIGH
OF HIGH











ACCUMULATION
ACCUMULATION
GROWTH













ACCUMULATION
ACCUMULATION

1⁄2t
1⁄2t















HOLDING
HOLDING
DEGREE OF
DEGREE OF
COOLING
ACCUMULATION
ACCUMULATION



TEMPERATURE
TIME
{200} PLANES
{222} PLANES
RATE
DEGREE OF
DEGREE OF


No.
° C.
sec
IN αFe
IN αFe
° C./sec
{200} PLANES
{222} PLANES





46
1100
10
15
14
50
15
14


47
1100
10
32
8.6
50
32
8.6


48
1100
10
52
1.9
50
52
1.9


49
1100
10
67
0.7
50
67
0.7


50
1100
10
71
0.3
50
71
0.3


51
1100
10
68
0.8
50
68
0.8


52
1100
10
44
3.8
50
44
3.8


53
1100
10
40
4.9
50
40
4.9


54
1250
30
16
13
150
16
13


55
1250
30
33
8.3
150
33
8.3


56
1250
30
51
1.9
150
51
1.9


57
1250
30
65
0.8
150
65
0.8


58
1250
30
70
0.3
150
70
0.3


59
1250
30
67
0.7
150
67
0.7


60
1250
30
43
3.9
150
43
3.9


61
1250
30
41
4.5
150
41
4.5


62
1050
100
16
12
20
16
12


63
1050
100
31
9.2
20
31
9.2


64
1050
100
50
2.4
20
50
2.4


65
1050
100
66
0.9
20
66
0.9


66
1050
100
69
0.4
20
69
0.4


67
1050
100
64
1.1
20
64
1.1


68
1050
100
42
4.1
20
42
4.1


69
1050
100
39
5.2
20
39
5.2


70
1150
200
14
14
5
14
14


71
1150
200
33
8.4
5
33
8.4


72
1150
200
53
1.5
5
53
1.5


73
1150
200
66
0.8
5
66
0.8


74
1150
200
70
0.4
5
70
0.4


75
1150
200
67
0.8
5
67
0.8


76
1150
200
44
3.4
5
44
3.4


77
1150
200
41
4.8
5
41
4.8


78
1300
500
17
12
250
17
12


79
1300
500
34
8.7
250
34
8.7


80
1300
500
54
1.4
250
54
1.4


81
1300
500
65
0.9
250
65
0.9


82
1300
500
68
0.6
250
68
0.6


83
1300
500
65
1
250
65
1


84
1300
500
43
3..5
250
43
3..5


85
1300
500
40
4.8
250
40
4.8













PRODUCT











TEXTURE EVALUATION












ACCUMULATION
ACCUMULATION

















DEGREE OF
DEGREE OF

MAGNETIC FLUX






{200} PLANES
{222} PLANES

DENSITY EVALUATION

α-REGION


















No.
IN αFe
IN αFe
Z
B50 T
Bs T
B50/Bs
ΔB T
NOTE
TEMPERATURE







46
15
14
0.5
1.59
2.02
0.79
0.090
COMPARATIVE
1050











EXAMPLE 11



47
32
8.6
2.5
1.73
2.02
0.86
0.080
PRESENT
950











INVENTION











EXAMPLE 36



48
52
1.9
12  
1.78
2.02
0.88
0.080
PRESENT
850











INVENTION











EXAMPLE 37



49
67
0.7
75  
1.89
2.02
0.94
0.030
PRESENT
750











INVENTION











EXAMPLE 38



50
71
0.3
143   
1.93
2.02
0.96
0.015
PRESENT
600











INVENTION











EXAMPLE 39



51
68
0.8
116   
1.92
2.02
0.95
0.019
PRESENT
450











INVENTION











EXAMPLE 40



52
44
3.8
2.9
1.76
2.02
0.87
0.115
PRESENT
300











INVENTION











EXAMPLE 41



53
40
4.9
0.9
1.75
2.02
0.87
0.200
COMPARATIVE
250











EXAMPLE 12



54
16
13
0.7
1.58
2.02
0.78
0.091
COMPARATIVE
1050











EXAMPLE 13



55
33
8.3
2.9
1.72
2.02
0.85
0.083
PRESENT
950











INVENTION











EXAMPLE 42



56
51
1.9
18  
1.79
2.02
0.89
0.070
PRESENT
850











INVENTION











EXAMPLE 43



57
65
0.8
83  
1.87
2.02
0.93
0.070
PRESENT
750











INVENTION











EXAMPLE 44



58
70
0.3
183   
1.93
2.02
0.96
0.045
PRESENT
600











INVENTION











EXAMPLE 45



59
67
0.7
127   
1.93
2.02
0.96
0.031
PRESENT
450











INVENTION











EXAMPLE 46



60
43
3.9
4.7
1.77
2.02
0.88
0.138
PRESENT
300











INVENTION











EXAMPLE 47



61
41
4.5
1.2
1.75
2.02
0.87
0.190
COMPARATIVE
250











EXAMPLE 14



62
16
12
1.1
1.59
2.02
0.79
0.090
COMPARATIVE
1050











EXAMPLE 15



63
31
9.2
2.4
1.73
2.02
0.86
0.080
PRESENT
950











INVENTION











EXAMPLE 48



64
50
2.4
15  
1.78
2.02
0.88
0.080
PRESENT
850











INVENTION











EXAMPLE 49



65
66
0.9
77  
1.87
2.02
0.93
0.076
PRESENT
750











INVENTION











EXAMPLE 50



66
69
0.4
125   
1.92
2.02
0.95
0.050
PRESENT
600











INVENTION











EXAMPLE 51



67
64
1.1
108   
1.92
2.02
0.95
0.042
PRESENT
450











INVENTION











EXAMPLE 52



68
42
4.1
2.6
1.77
2.02
0.88
0.138
PRESENT
300











INVENTION











EXAMPLE 53



69
39
5.2
1.4
1.76
2.02
0.87
0.220
COMPARATIVE
250











EXAMPLE 16



70
14
14
0.4
1.58
2.02
0.78
0.230
COMPARATIVE
1050











EXAMPLE 17



71
33
8.4
2.9
1.72
2.02
0.85
0.135
PRESENT
950











INVENTION











EXAMPLE 54



72
53
1.5
36  
1.77
2.02
0.88
0.094
PRESENT
850











INVENTION











EXAMPLE 55



73
66
0.8
98  
1.87
2.02
0.93
0.075
PRESENT
750











INVENTION











EXAMPLE 56



74
70
0.4
178   
1.94
2.02
0.96
0.061
PRESENT
600











INVENTION











EXAMPLE 57



75
67
0.8
47  
1.94
2.02
0.96
0.042
PRESENT
450











INVENTION











EXAMPLE 58



76
44
3.4
10.4 
1.76
2.02
0.87
0.137
PRESENT
300











INVENTION











EXAMPLE 59



77
41
4.8
1.2
1.75
2.02
0.87
0.230
COMPARATIVE
250











EXAMPLE 18



78
17
12
0.9
1.59
2.02
0.79
0.210
COMPARATIVE
1050











EXAMPLE 19



79
34
8.7
4.7
1.73
2.02
0.86
0.143
PRESENT
950











INVENTION











EXAMPLE 60



80
54
1.4
45  
1.78
2.02
0.88
0.090
PRESENT
850











INVENTION











EXAMPLE 61



81
65
0.9
118   
1.87
2.02
0.93
0.064
PRESENT
750











INVENTION











EXAMPLE 62



82
68
0.6
159   
1.92
2.02
0.95
0.020
PRESENT
600











INVENTION











EXAMPLE 63



83
65
1
69  
1.92
2.02
0.95
0.031
PRESENT
450











INVENTION











EXAMPLE 64



84
43
3..5
3.7
1.76
2.02
0.87
0.120
PRESENT
300











INVENTION











EXAMPLE 65



85
40
4.8
1.7
1.75
2.02
0.87
0.230
COMPARATIVE
250











EXAMPLE 20










In each of present invention examples, it was possible to confirm that Z is not less than 2.0 nor more than 200, the magnetic flux density difference ΔB becomes a small value as compared to comparative examples, and a high magnetic flux density can be obtained thoroughly in an in-plane circumferential direction. Further, in these Fe-based metal sheets, it was possible to confirm that an excellent magnetic property in which the value of B50/Bs is 0.80 or more is obtained.
Further, in the present invention examples, as shown in Table 2 to Table 4, it was possible to confirm that the (200) plane in the α-Fe phase is likely to be highly accumulated at each of the stages of the heat treatment.
Further, an L cross section of each of the present invention examples was observed, and thereby it was confirmed that the α single phase region made of the α single phase based component exists in at least a partial region including the surfaces and a ratio of the α single phase region to the L cross section is 1% or more.
When the Z value was not less than 2 nor more than 200 as defined in the present invention as above, it was possible to confirm that a high magnetic flux density is obtained thoroughly in the in-plane circumferential direction. Further, in order to obtain the Fe-based metal sheet as above, the α-region rolling was performed at a temperature of higher than 300° C. and lower than the A3 point between the hot rolling and the cold rolling, thereby making it possible to obtain an intended product.
In contrast to this, when the base metal sheets obtained by performing the rolling under the conditions not satisfying the requirements of the present invention were used, it was not possible to obtain a high magnetic flux density such as that in the present invention examples in the in-plane circumferential direction thoroughly.
Example 2
In this example, base metal sheets of No. 18 to No. 35 each made of a component C, D, or E shown in Table 1 were manufactured under various rolling conditions, to then have various different metals applied thereto as a second layer, and then Fe-based metal sheets were fabricated, of which the previously described Z value (=(A+0.97B)/0.98C) and the magnetic flux density difference ΔB were examined. Further, the relationship between various manufacturing conditions and an accumulation degree of {200} planes was also examined. Further, effects obtained by changing a starting temperature in an α-region rolling process were also examined in detail.
First, ingots each having the component C, D, or E shown in Table 1 and a balance being composed of Fe and inevitable impurities were melted by vacuum melting. Then, these were used as rolling materials to be worked into cold-rolled sheets (the base metal sheets) each having a predetermined thickness under conditions of hot rolling, α-region rolling, and cold rolling shown in Table 5 below.









TABLE 5









α-REGION










HOT ROLLING
ROLLING












START
FINISH
START
FINISH

















BASE


TEMPER-
THICK-
TEMPER-
THICK-
REDUC-
TEMPER-
THICK-
TEMPER-


MATERIAL

A3
ATURE
NESS
ATURE
NESS
TION
ATURE
NESS
ATURE


No.
COMPONENT
POINT
° C.
mm
° C.
mm
RATIO
° C.
mm
° C.





18
C
915
1050
200
930
60
−1.20
700
60
610


19
C
915
1050
200
930
30
−1.90
700
30
610


20
C
915
1050
200
930
20
−2.30
700
20
610


21
C
915
1050
200
930
10
−3.00
700
10
610


22
C
915
1050
200
930
8
−3.22
700
8
610


23
C
915
1050
200
930
4
−3.91
700
4
610


24
D
870
1050
300
930
15
−3.00
650
15
570


25
D
870
1050
150
930
15
−2.30
650
15
570


26
D
870
1050
75
930
15
−1.61
650
15
570


27
D
870
1050
50
930
15
−1.20
650
15
570


28
D
870
1050
20
930
15
−0.29
650
15
570


29
E
942
1200
240
1050
50
−1.67
750
30
670


30
E
942
1200
240
1050
50
−1.57
750
30
670


31
E
942
1200
240
1050
50
−1.57
750
30
670


32
E
942
1200
240
1050
50
−1.57
750
30
670


33
E
942
1200
240
1050
50
−1.57
750
30
670


34
E
942
1200
240
1050
50
−1.57
750
30
670


35
E
942
1200
240
1050
50
−1.57
750
30
670













α-REGION












ROLLING
COLD ROLLING














FINISH
START
FINISH
TOTAL

















BASE
THICK-
REDUC-
THICK-
THICK-
REDUC-
REDUC-
α REGION +



MATERIAL
NESS
TION
NESS
NESS
TION
TION
COLD



No.
mm
RATIO
mm
mm
RATIO
RATIO
ROLLING







18
2
−3.40
2
0.35
−1.74
−6.35
−5.14



19
2
−2.71
2
0.35
−1.74
−6.35
−4.45



20
2
−2.30
2
0.35
−1.74
−6.35
−4.05



21
2
−1.61
2
0.35
−1.74
−6.35
−3.35



22
2
−1.39
2
0.35
−1.74
−6.35
−3.13



23
2
−0.69
2
0.35
−1.74
−6.35
−2.44



24
3.5
−1.46
3.5
0.5
−1.95
−6.40
−3.40



25
3.5
−1.46
3.5
0.6
−1.95
−5.70
−3.40



26
3.5
−1.46
3.5
0.5
−1.95
−5.01
−3.40



21
3.5
−1.46
3.5
0.5
−1.95
−4.61
−3.40



28
3.5
−1.46
3.5
0.5
−1.95
−3.69
−3.40



29
6
−1.61
6
3
−0.69
−3.87
−2.30



30
6
−1.61
6
2
−1.10
−4.28
−2.71



31
6
−1.61
6
1
−1.79
−4.97
−3.40



32
6
−1.61
6
0.2
−3.40
−6.58
−5.01



33
6
−1.61
6
0.1
−4.09
−7.27
−5.70



34
6
−1.61
6
0.05
−4.79
−7.97
−6.40



35
6
−1.61
6
0.01
−6.40
−9.57
−8.01










In the case of the component C, first, the ingots each having a thickness of 200 mm heated to 1050° C. were each subjected to hot rolling at a reduction ratio of −1.20 to - 3.91 in terms of true strain, and hot-rolled sheets each having a thickness of 4 mm to 60 mm were obtained. Next, α-region rolling was started at 700° C., and these hot-rolled sheets were each subjected to the α-region rolling at a reduction ratio of −0.69 to −3.40 in terms of true strain to a thickness of 2 mm. Then, these rolled sheets were pickled, and then the base metal sheets were obtained by cold rolling. At this time, the reduction ratio was −1.74 in terms of true strain, and as a result, the thickness of each of the obtained base metal sheets was 0.35 mm.
In the case of the component D, first, the ingots each having a thickness of 20 mm to 300 mm heated to 1050° C. were each subjected to hot rolling at a reduction ratio of −0.29 to −3.00 in terms of true strain, and hot-rolled sheets each having a thickness of 15 mm were obtained. Next, α-region rolling was started at 650° C., and these hot-rolled sheets were each subjected to the α-region rolling at a reduction ratio of −1.46 in terms of true strain to a thickness of 3.5 mm. Then, these rolled sheets were pickled, and then the base metal sheets were obtained by cold rolling. At this time, the reduction ratio was −1.95 in terms of true strain, and as a result, the thickness of each of the obtained base metal sheets was 0.50 mm.
In the case of the component E, first, the ingots each having a thickness of 240 mm heated to 1200° C. were each subjected to hot rolling at a reduction ratio of −1.57 in terms of true strain, and hot-rolled sheets each having a thickness of 50 mm were obtained. Next, α-region rolling was started at 750° C., and these hot-rolled sheets were each subjected to the α-region rolling at a reduction ratio of −1.61 in terms of true strain to a thickness of 6.0 mm. Then, these rolled sheets were pickled, and then the base metal sheets were obtained by cold rolling. At this time, each of the reduction ratios was −0.69 to −6.40 in terms of true strain, and as a result, the thickness of each of the obtained base metal sheets was 0.01 mm to 3.0 mm.
With respect to the base metal sheets obtained by the above procedure, a texture of a surface layer portion of each of the base materials was measured by X-ray diffraction to obtain an accumulation degree of {200} planes and an accumulation degree of {222} planes by the previously described method. Further, thinning was performed so that a structure could be observed from a direction perpendicular to an L cross-section, and a region up to 1⁄4t from the surface was observed. The main phase of each of the obtained base metal sheets at room temperature was an α-Fe phase. Further, as a result of measurement, the A3 point at which the α-γtransformation occurred was 915° C. in the component C, 870° C. in the component D, and 942° C. in the component E.
Next, both surfaces of each of the base metal sheets of No. 18 to No. 35 shown in Table 5 were coated with each of various different metal elements as the second layer by a vapor deposition method, a sputtering method, an electroplating method, or a hot dipping method. As shown in Table 6 and Table 7 below, as the different metal element, any one of Al, Si, Ga, Sn, V, W, Mo, and Zn was selected. The thickness of each of the coatings was as shown in Table 6 and Table 7.
Next, an experiment was performed in which a heat treatment was performed on the base metal sheets to each of which the second layers were bonded under various conditions. As a method of the experiment, the experiment was performed by the same method described in Example 1. Further, the observation of a texture in this period was also performed by the same method described in Example 1.
Further, magnetometry was performed in the same manner as that in Example 1 in order to evaluate obtained products, and further the Z value was calculated by X-ray diffraction.
Table 6 and Table 7 below show the accumulation degrees of the {200} planes and the accumulation degrees of the {222} planes measured in the respective processes during the manufacture and after the manufacture, the Z values of the obtained Fe-based metal sheets, and evaluation results of the magnetometry.









TABLE 6









MANUFACTURE









SEEDING












BASE

SEEDING
SEEDING














MATERIAL
SECOND LAYER
TEMPERATURE

ACCUMULATION
ACCUMULATION


















THICK-

THICK-

INCREASING

DEGREE OF
DEGREE OF




NESS

NESS
BONDING
RATE
MEASURED
{200} PLANES
{222} PLANES


No.
No.
mm
ELEMENT
μm
METHOD
° C./sec
TEMPERATURE
IN αFe
IN αFe





86
18
0.35
Al
7
VAPOR
20
915
34
7.8







DEPOSITION


87
19
0.35
Al
7
VAPOR
20
915
34
7.9







DEPOSITION


88
20
0.35
Al
7
VAPOR
20
915
33
8.3







DEPOSITION


89
21
0.35
Al
7
VAPOR
20
915
30
9.3







DEPOSITION


90
22
0.35
Al
7
VAPOR
20
915
28
10







DEPOSITION


91
23
0.35
Al
7
VAPOR
20
915
26
10







DEPOSITION


92
18
0.35
Si
8
SPUTTERING
10
915
35
7.1


93
19
0.35
Si
8
SPUTTERING
10
915
34
7.3


94
20
0.35
Si
8
SPUTTERING
10
915
32
7.8


95
21
0.35
Si
8
SPUTTERING
10
915
27
9.7


96
22
0.35
Si
8
SPUTTERING
10
915
26
10


97
23
0.35
Si
8
SPUTTERING
10
915
25
11


98
18
0.35
Ga
6
VAPOR
0.5
915
32
8.5







DEPOSITION


99
19
0.35
Ga
6
VAPOR
0.5
915
32
8.6







DEPOSITION


100
20
0.35
Ga
6
VAPOR
0.5
915
31
8.9







DEPOSITION


101
21
0.35
Ga
6
VAPOR
0.5
915
28
9.8







DEPOSITION


102
22
0.35
Ga
6
VAPOR
0.5
915
26
10







DEPOSITION


103
23
0.35
Ga
6
VAPOR
0.5
915
25
11







DEPOSITION


104
18
0.35
Sn
10
ELECTROLYTIC
5
915
36
6.5







PLATING


105
19
0.35
Sn
10
ELECTROLYTIC
5
915
35
6.7







PLATING


106
20
0.35
Sn
10
ELECTROLYTIC
5
915
36
7.2







PLATING


107
21
0.35
Sn
10
ELECTROLYTIC
5
915
31
8.7







PLATING


108
22
0.35
Sn
10
ELECTROLYTIC
5
915
27
9.3







PLATING


109
23
0.35
Sn
10
ELECTROLYTIC
5
915
25
11







PLATING


110
18
0.35
V
11
SPUTTERING
10
915
34
7.9


111
19
0.35
V
11
SPUTTERING
10
915
33
8.2


112
20
0.35
V
11
SPUTTERING
10
915
31
8.6


113
21
0.35
V
11
SPUTTERING
10
915
28
9.7


114
22
0.35
V
11
SPUTTERING
10
915
27
10


115
23
0.35
V
11
SPUTTERING
10
915
25
10


116
18
0.35
W
6
SPUTTERING
0.5
915
34
7.6


117
19
0.35
W
6
SPUTTERING
0.5
915
33
8.2


118
20
0.35
W
6
SPUTTERING
0.5
915
31
9.2


119
21
0.35
W
6
SPUTTERING
0.5
915
28
10


120
22
0.35
W
6
SPUTTERING
0.5
915
27
10


121
23
0.35
W
6
SPUTTERING
0.5
915
25
11












MANUFACTURE










SAVING AND ACHIEVEMENT OF




HIGH ACCUMULATION
















ACHIEVEMENT
ACHIEVEMENT






OF HIGH
OF HIGH











ACCUMULATION
ACCUMULATION
GROWTH













ACCUMULATION
ACCUMULATION

1⁄2t
1⁄2t















HOLDING
HOLDING
DEGREE OF
DEGREE OF
COOLING
ACCUMULATION
ACCUMULATION



TEMPERATURE
TIME
{200} PLANES
{222} PLANES
RATE
DEGREE OF
DEGREE OF


No.
° C.
sec
IN αFe
IN αFe
° C./sec
{200} PLANES
{222} PLANES





86
1000
40
63
0.8
100
63
0.8


87
1000
40
62
0.8
100
62
0.8


88
1000
40
61
0.9
100
61
0.9


89
1000
40
50
3.9
100
50
3.9


90
1000
40
42
5.7
100
42
5.7


91
1000
40
32
8.5
100
32
8.5


92
1050
25
62
0.9
20
62
0.9


93
1050
25
61
0.9
20
61
0.9


94
1050
25
60
1.2
20
60
1.2


95
1050
25
45
3.5
20
45
3.5


96
1050
25
42
5.3
20
42
5.3


97
1050
25
31
9.1
20
31
9.1


98
950
120
60
1.1
50
60
1.1


99
950
120
60
1
50
60
1


100
950
120
59
1.2
50
59
1.2


101
950
120
43
4.5
50
43
4.5


102
950
120
41
6.2
50
41
6.2


103
950
120
30
9.7
50
30
9.7


104
1000
10
64
0.7
200
64
0.7


105
1000
10
64
0.7
200
64
0.7


106
1000
10
63
0.8
200
63
0.8


107
1000
10
50
2.8
200
50
2.8


108
1000
10
43
4.7
200
43
4.7


109
1000
10
33
7.9
200
33
7.9


110
1200
15
61
0.9
250
61
0.9


111
1200
15
61
0.9
250
61
0.9


112
1200
15
60
1.2
250
60
1.2


113
1200
15
45
4.3
250
45
4.3


114
1200
15
39
6.6
250
39
6.6


115
1200
15
30
8.2
250
30
8.2


116
1300
30
60
1.1
80
60
1.1


117
1300
30
59
1.3
80
59
1.3


118
1300
30
58
1.6
80
58
1.6


119
1300
30
46
3.5
80
46
3.5


120
1300
30
38
6.9
80
38
8.9


121
1300
30
31
9.1
80
31
9.1













PRODUCT











TEXTURE EVALUATION












ACCUMULATION
ACCUMULATION

















DEGREE OF
DEGREE OF

MAGNETIC FLUX






{200} PLANES
{222} PLANES

DENSITY EVALUATION

α-REGION


















No.
IN αFe
IN αFe
Z
B50 T
Bs T
B50/Bs
ΔB T
NOTE
TEMPERATURE







86
63
0.8
135
1.87
2.04
0.92
0.038
PRESENT
700











INVENTION











EXAMPLE 66



87
62
0.8
120
1.86
2.04
0.91
0.048
PRESENT
700











INVENTION











EXAMPLE 67



88
61
0.9
52
1.86
2.04
0.91
0.053
PRESENT
700











INVENTION











EXAMPLE 68



89
47
3.9
24
1.82
2.04
0.89
0.068
PRESENT
700











INVENTION











EXAMPLE 89



90
40
5.7
8.6
1.76
2.04
0.86
0.076
PRESENT
700











INVENTION











EXAMPLE 70



91
32
8.5
2.5
1.72
2.04
0.84
0.091
PRESENT
700











INVENTION











EXAMPLE 71



92
62
0.9
157
1.86
2.04
0.91
0.035
PRESENT
700











INVENTION











EXAMPLE 72



93
61
0.9
132
1.86
2.04
0.91
0.046
PRESENT
700











INVENTION











EXAMPLE 73



94
60
1.2
62
1.85
2.04
0.91
0.057
PRESENT
700











INVENTION











EXAMPLE 74



95
45
3.5
28
1.78
2.04
0.87
0.066
PRESENT
700











INVENTION











EXAMPLE 75



96
42
5.3
9.4
1.76
2.04
0.86
0.084
PRESENT
700











INVENTION











EXAMPLE 76



97
31
9.1
3.8
1.73
2.04
0.85
0.091
PRESENT
700











INVENTION











EXAMPLE 77



98
60
1.1
167
1.86
2.04
0.91
0.031
PRESENT
700











INVENTION











EXAMPLE 78



99
60
1
121
1.86
2.04
0.91
0.047
PRESENT
700











INVENTION











EXAMPLE 79



100
59
1.2
71
1.85
2.04
0.91
0.054
PRESENT
700











INVENTION











EXAMPLE 80



101
43
4.5
31
1.79
2.04
0.88
0.068
PRESENT
700











INVENTION











EXAMPLE 81



102
41
6.2
10
1.76
2.04
0.86
0.079
PRESENT
700











INVENTION











EXAMPLE 82



103
30
9.7
2.6
1.72
2.04
0.84
0.087
PRESENT
700











INVENTION











EXAMPLE 83



104
64
0.7
184
1.88
2.04
0.92
0.027
PRESENT
700











INVENTION











EXAMPLE 84



105
64
0.7
137
1.87
2.04
0.92
0.044
PRESENT
700











INVENTION











EXAMPLE 85



106
63
0.8
68
1.88
2.04
0.92
0.057
PRESENT
700











INVENTION











EXAMPLE 86



107
50
2.8
32
1.81
2.04
0.89
0.071
PRESENT
700











INVENTION











EXAMPLE 87



108
43
4.7
9.4
1.75
2.04
0.86
0.082
PRESENT
700











INVENTION











EXAMPLE 88



109
33
7.9
3.1
1.71
2.04
0.84
0.094
PRESENT
700











INVENTION











EXAMPLE 89



110
61
0.9
154
1.87
2.04
0.92
0.022
PRESENT
700











INVENTION











EXAMPLE 90



111
61
0.9
118
1.87
2.04
0.92
0.039
PRESENT
700











INVENTION











EXAMPLE 91



112
60
1.2
66
1.86
2.04
0.91
0.053
PRESENT
700











INVENTION











EXAMPLE 92



113
45
4.3
24
1.81
2.04
0.89
0.067
PRESENT
700











INVENTION











EXAMPLE 93



114
39
6.6
8.9
1.76
2.04
0.86
0.075
PRESENT
700











INVENTION











EXAMPLE 94



115
30
8.2
4.2
1.72
2.04
0.84
0.088
PRESENT
700











INVENTION











EXAMPLE 95



116
60
1.1
186
1.86
2.04
0.91
0.019
PRESENT
700











INVENTION











EXAMPLE 96



117
59
1.3
136
1.85
2.04
0.91
0.032
PRESENT
700











INVENTION











EXAMPLE 97



118
58
1.6
74
1.85
2.04
0.91
0.050
PRESENT
700











INVENTION











EXAMPLE 98



119
46
3.5
28
1.79
2.04
0.88
0.062
PRESENT
700











INVENTION











EXAMPLE 99



120
38
6.9
12
1.76
2.04
0.86
0.072
PRESENT
700











INVENTION











EXAMPLE 100



121
31
9.1
3.9
1.71
2.04
0.84
0.093
PRESENT
700











INVENTION











EXAMPLE 101



















TABLE 7









MANUFACTURE









SEEDING












BASE

SEEDING
SEEDING














MATERIAL
SECOND LAYER
TEMPERATURE

ACCUMULATION
ACCUMULATION


















THICK-

THICK-

INCREASING

DEGREE OF
DEGREE OF




NESS

NESS
BONDING
RATE
MEASURED
{200} PLANES
{222} PLANES


No.
No.
mm
ELEMENT
μm
METHOD
° C./sec
TEMPERATURE
IN αFe
IN αFe





122
24
0.5
Al
10
VAPOR
10
870
31
7.7







DEPOSITION


123
25
0.5
Al
10
VAPOR
10
870
31
7.9







DEPOSITION


124
26
0.5
Al
10
VAPOR
10
870
30
9.3







DEPOSITION


125
27
0.5
Al
10
VAPOR
10
870
27
9.8







DEPOSITION


126
28
0.5
Al
10
VAPOR
10
870
25
10







DEPOSITION


127
24
0.5
Si
12
VAPOR
20
870
31
8.1







DEPOSITION


128
25
0.5
Si
12
VAPOR
20
870
31
8.2







DEPOSITION


129
26
0.5
Si
12
VAPOR
20
870
30
9.3







DEPOSITION


130
27
0.5
Si
12
VAPOR
20
870
27
10







DEPOSITION


131
28
0.5
Si
12
VAPOR
20
870
26
11







DEPOSITION


132
24
0.5
Mo
8
SPUTTERING
1
870
33
6.8


133
25
0.5
Mo
8
SPUTTERING
1
870
32
7.3


134
26
0.5
Mo
8
SPUTTERING
1
870
30
8.8


135
27
0.5
Mo
8
SPUTTERING
1
870
27
9.3


136
28
0.5
Mo
8
SPUTTERING
1
870
25
10


137
29
3
Al
120
HOT DIPPING
2
942
13
13


138
30
2
Al
80
HOT DIPPING
2
942
25
10


139
31
1
Al
40
HOT DIPPING
2
942
31
8.3


140
32
0.2
Al
8
VAPOR
2
942
32
7.5







DEPOSITION


141
33
0.1
Al
4
VAPOR
2
942
33
6.7







DEPOSITION


142
34
0.05
Al
2
VAPOR
2
942
33
6.5







DEPOSITION


143
35
0.01
Al
0.4
VAPOR
2
942
32
6.4







DEPOSITION


144
29
3
Sn
60
HOT DIPPING
5
942
12
12


145
30
2
Sn
40
HOT DIPPING
5
942
25
10


146
31
1
Sn
20
HOT DIPPING
5
942
32
8.1


147
32
0.2
Sn
4
ELECTROLYTIC
5
942
33
7.1







PLATING


148
33
0.1
Sn
2
ELECTROLYTIC
5
942
34
6.3







PLATING


149
34
0.05
Sn
1
ELECTROLYTIC
5
942
35
6.1







PLATING


150
35
0.01
Sn
0.2
ELECTROLYTIC
5
942
34
6.6







PLATING


151
29
3
Zn
60
HOT DIPPING
1
942
14
13


152
30
2
Zn
40
HOT DIPPING
1
942
25
11


153
31
1
Zn
20
HOT DIPPING
1
942
30
8.8


154
32
0.2
Zn
4
ELECTROLYTIC
1
942
31
7.8







PLATING


155
33
0.1
Zn
2
ELECTROLYTIC
1
942
32
6.5







PLATING


156
34
0.05
Zn
1
ELECTROLYTIC
1
942
32
6.3







PLATING


157
35
0.01
Zn
0.2
ELECTROLYTIC
1
942
32
6.7







PLATING












MANUFACTURE










SAVING AND ACHIEVEMENT OF




HIGH ACCUMULATION
















ACHIEVEMENT
ACHIEVEMENT






OF HIGH
OF HIGH











ACCUMULATION
ACCUMULATION
GROWTH













ACCUMULATION
ACCUMULATION

1⁄2t
1⁄2t















HOLDING
HOLDING
DEGREE OF
DEGREE OF
COOLING
ACCUMULATION
ACCUMULATION



TEMPERATURE
TIME
{200} PLANES
{222} PLANES
RATE
DEGREE OF
DEGREE OF


No.
° C.
sec
IN αFe
IN αFe
° C./sec
{200} PLANES
{222} PLANES





122
930
20
56
1.6
80
56
1.6


123
930
20
55
1.8
80
55
1.8


124
930
20
52
2.5
80
52
2.5


125
930
20
40
5.9
80
40
5.9


126
930
20
32
9.3
80
32
9.3


127
980
60
54
1.7
20
54
1.7


128
980
60
53
1.9
20
53
1.9


129
980
60
51
2.8
20
51
2.8


130
980
60
39
7.4
20
39
7.4


131
980
60
33
9.5
20
33
9.5


132
1000
15
56
1.7
50
56
1.7


133
1000
15
56
1.7
50
56
1.7


134
1000
15
53
2.1
50
53
2.1


135
1000
15
41
6.3
50
41
6.3


136
1000
15
31
9.3
50
31
9.3


137
1050
25
15
13
100
15
13


138
1050
25
32
9.5
100
32
9.5


139
1050
25
50
2.8
100
50
2.8


140
1050
25
54
2.1
100
54
2.1


141
1050
25
55
1.8
100
55
1.8


142
1050
25
56
1.7
100
56
1.7


143
1050
25
55
1.8
100
55
1.8


144
1100
60
14
14
200
14
14


145
1100
60
32
9.4
200
32
9.4


146
1100
60
51
2.5
200
51
2.5


147
1100
60
56
1.8
200
56
1.8


148
1100
60
57
1.3
200
57
1.3


149
1100
60
57
1.1
200
57
1.1


150
1100
60
56
1.4
200
56
1.4


151
980
200
16
12
50
16
12


152
980
200
30
9.8
50
30
9.8


153
980
200
50
3.1
50
50
3.1


154
980
200
52
2.5
50
52
2.5


155
980
200
54
2.1
50
54
2.1


156
980
200
55
1.9
50
55
1.9


157
980
200
54
2.1
50
54
2.1













PRODUCT











TEXTURE EVALUATION












ACCUMULATION
ACCUMULATION

















DEGREE OF
DEGREE OF

MAGNETIC FLUX






{200} PLANES
{222} PLANES

DENSITY EVALUATION

α-REGION


















No.
IN αFe
IN αFe
Z
B50 T
Bs T
B50/Bs
ΔB T
NOTE
TEMPERATURE







122
56
1.6
98
1.85
1.98
0.93
0.024
PRESENT
650











INVENTION











EXAMPLE 102



123
55
1.8
78
1.85
1.98
0.93
0.028
PRESENT
650











INVENTION











EXAMPLE 103



124
52
2.5
57
1.83
1.98
0.92
0.041
PRESENT
650











INVENTION











EXAMPLE 104



125
40
5.9
24
1.73
1.98
0.87
0.057
PRESENT
650











INVENTION











EXAMPLE 105



126
32
9.3
  3.8
1.69
1.98
0.85
0.087
PRESENT
650











INVENTION











EXAMPLE 106



127
54
1.7
110 
1.84
1.98
0.93
0.021
PRESENT
650











INVENTION











EXAMPLE 107



128
53
1.9
76
1.84
1.98
0.93
0.025
PRESENT
650











INVENTION











EXAMPLE 108



129
51
2.8
65
1.82
1.98
0.92
0.035
PRESENT
650











INVENTION











EXAMPLE 109



130
39
7.4
29
1.76
1.98
0.89
0.053
PRESENT
650











INVENTION











EXAMPLE 110



131
33
9.5
  5.6
1.68
1.98
0.85
0.086
PRESENT
650











INVENTION











EXAMPLE 111



132
56
1.7
105 
1.86
1.98
0.94
0.023
PRESENT
650











INVENTION











EXAMPLE 112



133
56
1.7
68
1.86
1.98
0.94
0.031
PRESENT
650











INVENTION











EXAMPLE 113



134
53
2.1
59
1.84
1.98
0.93
0.045
PRESENT
650











INVENTION











EXAMPLE 114



135
41
6.3
23
1.77
1.98
0.89
0.072
PRESENT
650











INVENTION











EXAMPLE 115



136
31
9.3
  4.2
1.68
1.98
0.85
0.092
PRESENT
650











INVENTION











EXAMPLE 116



137
15
13
  0.9
1.59
2.02
0.79
0.086
COMPARATIVE
750











EXAMPLE 21



138
32
9.5
  2.5
1.73
2.02
0.86
0.062
PRESENT
750











INVENTION











EXAMPLE 117



139
50
2.8
35
1.79
2.02
0.89
0.053
PRESENT
750











INVENTION











EXAMPLE 118



140
54
2.1
65
1.83
2.02
0.91
0.041
PRESENT
750











INVENTION











EXAMPLE 119



141
55
1.8
114 
1.83
2.02
0.91
0.032
PRESENT
750











INVENTION











EXAMPLE 120



142
56
1.7
126 
1.83
2.02
0.91
0.018
PRESENT
750











INVENTION











EXAMPLE 121



143
55
1.8
132 
1.83
2.02
0.91
0.015
PRESENT
750











INVENTION











EXAMPLE 122



144
14
14
  0.4
1.60
2.02
0.79
0.092
COMPARATIVE
750











EXAMPLE 22



145
32
9.4
  3.2
1.72
2.02
0.85
0.068
PRESENT
750











INVENTION











EXAMPLE 123



146
51
2.5
29
1.76
2.02
0.87
0.052
PRESENT
750











INVENTION











EXAMPLE 124



147
56
1.8
59
1.84
2.02
0.91
0.043
PRESENT
750











INVENTION











EXAMPLE 125



148
57
1.3
94
1.84
2.02
0.91
0.029
PRESENT
750











INVENTION











EXAMPLE 126



149
57
1.1
123 
1.85
2.02
0.92
0.021
PRESENT
750











INVENTION











EXAMPLE 127



150
56
1.4
135 
1.84
2.02
0.91
0.018
PRESENT
750











INVENTION











EXAMPLE 128



151
16
12
  1.1
1.58
2.02
0.78
0.087
COMPARATIVE
750











EXAMPLE 23



152
30
9.8
  4.5
1.71
2.02
0.85
0.058
PRESENT
750











INVENTION











EXAMPLE 129



153
50
3.1
27
1.79
2.02
0.89
0.047
PRESENT
750











INVENTION











EXAMPLE 130



154
52
2.5
49
1.83
2.02
0.91
0.039
PRESENT
750











INVENTION











EXAMPLE 131



155
54
2.1
79
1.83
2.02
0.91
0.025
PRESENT
750











INVENTION











EXAMPLE 132



156
55
1.9
132 
1.83
2.02
0.91
0.018
PRESENT
750











INVENTION











EXAMPLE 133



157
54
2.1
172 
1.83
2.02
0.91
0.012
PRESENT
750











INVENTION











EXAMPLE 134










In each of present invention examples, it was possible to confirm that the magnetic flux density difference ΔB becomes a small value as compared to comparative examples, and a high magnetic flux density is obtained thoroughly in an in-plane circumferential direction. Further, in these Fe-based metal sheets, it was possible to confirm that an excellent magnetic property in which the value of B50/Bs is 0.86 or more is obtained.
Further, in the present invention examples, as shown in Table 5 to Table 7, it was possible to confirm that the {200} plane in the α-Fe phase is likely to be highly accumulated at each of the stages of the heat treatment.
Further, an L cross section of each of the present invention examples was observed, and thereby it was confirmed that the α single phase region made of the α single phase based component exists in at least a partial region including the surfaces and a ratio of the α single phase region to the L cross section is 1% or more.
When the Z value was not less than 2 nor more than 200 as defined in the present invention as above, it was possible to confirm that a high magnetic flux density is obtained thoroughly in the in-plane circumferential direction. Further, in order to obtain the Fe-based metal sheet as above, the α-region rolling was performed at a temperature of 300° C. or higher and lower than the A3 point between the hot rolling and the cold rolling, thereby making it possible to obtain an intended product.
In contrast to this, when the base metal sheets obtained by performing the α-region rolling under the condition not satisfying the requirements of the present invention were used, it was not possible to obtain a high magnetic flux density such as that in the present invention examples in the in-plane circumferential direction thoroughly.
Example 3
In this example, as base metal sheets, Fe-based metal sheets were fabricated in a manner that pure irons each containing C: 0.050 mass %, Si: 0.0001 mass %, and Al: 0.0002 mass %, and having a balance being composed of Fe and inevitable impurities were subjected to decarburization to have Al applied thereto as a second layer, of which the previously described Z value (=(A+0.97B)/0.98C) and the magnetic flux density difference ΔB were examined. Further, the relationship between manufacturing conditions and an accumulation degree of {200} planes was also examined.
First, ingots were melted by vacuum melting, and then were subjected to hot rolling and cold rolling to be worked to a predetermined thickness, and the base metal sheets each composed of the previously described composition were obtained. Incidentally, the A1 point of the base metal sheets was 727° C.
In the hot rolling, the ingots each having a thickness of 230 mm heated to 1000° C. were thinned down to a thickness of 50 mm, and hot-rolled sheets were obtained. Sheet materials having various thicknesses were cut out from these hot-rolled sheets by machining and then were subjected to the cold rolling, and thereby cold-rolled sheets each having a thickness of 8 μm to 750 μm (the base metal sheets) were obtained.
Incidentally, the main phase of each of the base metal sheets at room temperature was an α-Fe phase and as a result of measurement, the A3 point at which the α-γ transformation occurred was 911° C. Further, a texture in the α-Fe phase of each of the base metal sheets was measured by X-ray diffraction, and by the previously described method, an accumulation degree of {200} planes and an accumulation degree of {222} planes were obtained. Further, as a result that up to the cold rolling was performed, it was confirmed that of each of the base metal sheets, the accumulation degree of the {200} planes is 20 to 26% and the accumulation degree of the {222} planes is 18 to 24%.
Next, these base metal sheets were subjected to decarburization annealing so that a decarburized depth (a distance x) might become 1 μm to 59 μm. A decarburization condition was set that the temperature is 800° C. and the decarburization time is 0.05 minutes to 550 minutes. With regard to the atmosphere during the decarburization annealing, a strong decarburizing atmosphere was applied in the case of the decarburization annealing being performed for one minute or shorter, and in the case of the decarburization annealing being performed for longer than one minute, a weak decarburizing atmosphere was applied in the first half of the decarburization annealing and a strong decarburizing atmosphere was applied in the second half of the decarburization annealing.
Then, after the decarburization annealing was performed, the decarburized depth and the C content of a decarburized region were measured and a structure and a crystal orientation of a surface layer were examined. The measurement of the crystal orientation was performed by the method using the previously described X-ray diffraction method, the X-ray was emitted from the surface, and the accumulation degree of the {200} planes in the α-Fe phase was obtained.
After the decarburization annealing, both surfaces of each of the base metal sheets were coated with Al as the second layer by an ion plating method (hereinafter, an IP method) to each have a thickness of 1 μm.
Next, an experiment was performed in which a heat treatment was performed on the base metal sheets to each of which the second layers was bonded under various conditions. A gold image furnace was used for the heat treatment, and a temperature increasing rate, a holding temperature, and a holding time were variously controlled by program control. During which the temperature increased to be held, the heat treatment was performed in an atmosphere vacuumed to a pressure of 10−3 Pa level. At the time of cooling, in the case of a cooling rate of 1° C./sec or lower, temperature control was performed in a vacuum by furnace output control. Further, in the case of the cooling rate of 10° C./sec or more, an Ar gas was introduced and the cooling rate was controlled by adjustment of its flow rate.
Further, the observation of the texture in this period was also performed by the same method described in Example 1. Further, magnetometry was performed in the same manner as that in Example 1 in order to evaluate obtained products, and further the Z value was calculated by the X-ray diffraction.
Further, an alloyed ratio in the second layer and a ratio of the α single phase region were defined and obtained as follows.
Plane distribution of the Fe content and plane distribution of the Al content were measured by using an EPMA (Electron Probe Micro-Analysis) method, with a field of view of an L direction 1 mm×the total thickness in an L cross section. First, as the alloyed ratio in the second layer, areas of a region satisfying Fe≦mass % and Al≧99.5 mass % before and after the heat treatment were obtained. Then, the alloyed ratio of the second layer was defined as (S0−S)/S0×100, where an area when Al was applied and the heat treatment was not performed was set to S0 and an area in the Fe-based metal sheet on which the whole heat treatment was completed was set to S.
Further, the ratio of the α single phase region was defined as (T/T0)×100, where an area of a cross section of the Fe-based metal sheet after the heat treatment, observed in the L cross section was set to T0 and an area of a diffused region of the different metal after the heat treatment was set to T. Incidentally, when the second layer was Al, an area of a region satisfying Al≧0.9 mass % was set to T.
Table 8 shows the base metal sheets and conditions of the decarburization and the heat treatment, and shows the accumulation degrees of the {200} planes and the accumulation degrees of the {200} planes measured during the manufacture (after the decarburization annealing) and after the manufacture (after the diffusion treatment), the Z values of the obtained Fe-based metal sheets, the alloyed ratios of the second layers, and evaluation results of the magnetometry.














TABLE 8














C CONTENT



BASE


DECARBU-
DECARBU-
DECARBU-
AFTER



MATERIAL
SHEET
DECARBU-
RIZATION
RIZATION
RIZED
DECARBU-



C CONTENT
THICKNESS
RIZING
TEMPERATURE
TIME
REGION
RIZATION


No.
mass %
μm
ATMOSPHERE
° C.
MINUTE
μm
mass %





201
0.050
10
STRONG
800
1
9
0.010


202
0.050
100
WEAK + STRONG
800
3
12
0.011


203
0.050
250
WEAK + STRONG
800
5
14
0.015


204
0.050
500
WEAK + STRONG
800
15
22
0.018


205
0.050
750
WEAK + STRONG
800
30
31
0.018


206
0.050
100
STRONG
800
0.1
6
0.008


207
0.050
100
WEAK + STRONG
800
250
36
0.017


208
0.050
500
WEAK + STRONG
800
550
49
0.017


209
0.050
200
WEAK + STRONG
800
10
18
0.008


210
0.050
200
WEAK + STRONG
800
10
18
0.009


211
0.050
200
WEAK + STRONG
800
10
19
0.008


212
0.050
200
WEAK + STRONG
800
10
18
0.010


213
0.050
200
WEAK + STRONG
800
10
17
0.008


214
0.050
200
WEAK + STRONG
800
10
18
0.009


215
0.050
200
WEAK + STRONG
800
10
18
0.009


216
0.050
200
WEAK + STRONG
800
10
16
0.010


217
0.050
150
WEAK + STRONG
800
8
15
0.007


218
0.050
150
WEAK + STRONG
800
8
14
0.006


219
0.050
150
WEAK + STRONG
800
8
16
0.007


220
0.050
150
WEAK + STRONG
800
8
14
0.007


221
0.050
150
WEAK + STRONG
800
8
14
0.007


222
0.050
150
WEAK + STRONG
800
8
15
0.007


223
0.050
150
WEAK + STRONG
800
8
16
0.006


224
0.050
300
WEAK + STRONG
800
15
21
0.011


225
0.050
300
WEAK + STRONG
800
15
22
0.009


226
0.050
300
WEAK + STRONG
800
15
22
0.009


227
0.050
8
STRONG
800
1
8
0.010


228
0.050
100
STRONG
800
0.05
1
0.050


229
0.050
100
WEAK + STRONG
800
60
59
0.003


230
0.050
100
WEAK + STRONG
800
18
23
0.010


231
0.050
100
WEAK + STRONG
800
18
25
0.011


232
0.050
100
WEAK + STRONG
800
18
26
0.009


233
0.050
100
WEAK + STRONG
800
18
25
0.009


234
0.050
100
WEAK + STRONG
800
18
24
0.010


235
0.050
100
WEAK + STRONG
800
18
26
0.009


















ACCUMULATION









DEGREE

TEMPERATURE



OF {200}
FERRITE-
INCREASING
HOLDING
HOLDING
COOLING



PLANES AFTER
FORMING
RATE
TEMPERATURE
TIME
RATE
(S0 − S)/


No.
DECARBURIZATION
ELEMENT
° C./sec
T1 ° C.
MINUTE
° C./sec
S0 × 100





201
26
Al
0.5
1000
5
100
79


202
24
Al
0.5
1000
5
100
65


203
25
Al
0.5
1000
5
100
52


204
21
Al
0.5
1000
5
100
39


205
29
Al
0.5
1000
5
100
37


206
23
Al
0.5
1000
5
100
66


207
26
Al
0.5
1000
5
100
64


208
38
Al
0.5
1000
5
100
31


209
26
Al
0.1
950
1
100
61


210
27
Al
1
1000
1
100
59


211
26
Al
5
1000
5
100
62


212
25
Al
10
1000
5
100
60


213
26
Al
20
1000
5
100
58


214
26
Al
0.5
950
5
100
59


215
27
Al
0.5
1050
5
100
60


216
25
Al
0.5
1200
5
100
57


217
28
Al
0.5
1000
0.5
100
29


218
29
Al
0.5
1000
10
100
61


219
30
Al
0.5
1000
30
100
76


220
29
Al
0.5
1000
60
100
81


221
29
Al
0.5
1000
120
100
96


222
28
Al
0.5
1000
550
100
100


223
30
Al
0.5
1000
4500
100
100


224
22
Al
0.5
1000
10
0.1
79


225
21
Al
0.5
1000
10
10
51


226
22
Al
0.5
1000
10
450
55


227
26
Al
0.5
950
1
100
100


228
17
Al
0.5
1000
10
100
65


229
28
Al
0.5
1000
10
100
62


230
26
NONE
0.5
1000
10
100
0


231
24
Al
0.5
900
10
100
46


232
27
Al
0.5
1350
10
100
78


233
23
Al
0.5
1000
6000
100
100


234
25
Al
0.5
1000
10
0.05
85


235
28
Al
0.5
1000
10
550
34





















ACCUMULATION
ACCUMULATION









DEGREE OF
DEGREE OF





{200} PLANES
{222} PLANES

B50/Bs




T/T0 ×
AFTER
AFTER

OF
ΔB



No.
100
DIFFUSION
DIFFUSION
Z
PRODUCT
T
NOTE







201
64
54
16
124 
0.892
0.042
INVENTION










EXAMPLE 201



202
50
42
28
56
0.864
0.057
INVENTION










EXAMPLE 202



203
43
36
24
  8.9
0.842
0.098
INVENTION










EXAMPLE 203



204
32
36
20
11
0.852
0.091
INVENTION










EXAMPLE 204



205
30
37
22
15
0.859
0.085
INVENTION










EXAMPLE 205



206
49
46
24
69
0.893
0.042
INVENTION










EXAMPLE 206



207
44
41
22
42
0.865
0.054
INVENTION










EXAMPLE 207



208
26
32
27
  3.1
0.833
0.101
INVENTION










EXAMPLE 208



209
50
41
23
39
0.859
0.083
INVENTION










EXAMPLE 209



210
48
39
26
25
0.865
0.071
INVENTION










EXAMPLE 210



211
51
42
18
68
0.872
0.045
INVENTION










EXAMPLE 211



212
46
34
25
  4.3
0.851
0.096
INVENTION










EXAMPLE 212



213
50
49
11
76
0.897
0.038
INVENTION










EXAMPLE 213



214
48
36
23
16
0.845
0.084
INVENTION










EXAMPLE 214



215
44
48
18
82
0.896
0.021
INVENTION










EXAMPLE 215



216
46
60
9
148 
0.904
0.016
INVENTION










EXAMPLE 216



217
24
34
24
  3.5
0.835
0.115
INVENTION










EXAMPLE 217



218
47
73
6
175 
0.953
0.008
INVENTION










EXAMPLE 218



219
55
62
11
152 
0.913
0.011
INVENTION










EXAMPLE 219



220
68
58
14
135 
0.901
0.018
INVENTION










EXAMPLE 220



221
75
52
15
112 
0.899
0.021
INVENTION










EXAMPLE 221



222
74
59
10
139 
0.908
0.016
INVENTION










EXAMPLE 222



223
76
55
12
131 
0.895
0.018
INVENTION










EXAMPLE 223



224
64
63
8
162 
0.918
0.011
INVENTION










EXAMPLE 224



225
43
72
4
189 
0.958
0.007
INVENTION










EXAMPLE 225



226
37
68
6
158 
0.954
0.009
INVENTION










EXAMPLE 226



227
100
18
31
  0.3
0.789
0.108
COMPARATIVE










EXAMPLE 201



228
41
24
37
  1.2
0.785
0.110
COMPARATIVE










EXAMPLE 202



229
38
37
24
19
0.857
0.086
INVENTION










EXAMPLE 227



230
0
19
31
  1.4
0.778
0.105
COMPARATIVE










EXAMPLE 203



231
37
22
27
  0.9
0.768
0.104
COMPARATIVE










EXAMPLE 204



232
51
38
20
14
0.832
0.095
INVENTION










EXAMPLE 228



233
76
38
26
17
0.842
0.091
INVENTION










EXAMPLE 229



234
70
37
25
16
0.845
0.089
INVENTION










EXAMPLE 230



235
27
38
26
15
0.841
0.089
INVENTION










EXAMPLE 231










As shown in Table 8, in each of present invention examples, it was possible to confirm that the magnetic flux density difference ΔB becomes a small value as compared to comparative examples and a high magnetic flux density can be obtained thoroughly in the in-plane circumferential direction. Further, in these Fe-based metal sheets, it was possible to confirm that an excellent magnetic property in which the value of B50/Bs is 0.80 or more is obtained.
Further, in the present invention examples, it was possible to confirm that the alloyed ratio and the ratio of the α single phase region can be controlled by the combination of the decarburized depth of the base metal sheet, the temperature increasing rate, the holding temperature after the heating, and the holding time, and the Fe-based metal sheet having an excellent magnetic property can be obtained.
Further, an L cross section of each of the present invention examples was observed, and thereby it was confirmed that the α single phase region made of the α single phase based component exists in at least a partial region including the surfaces and the ratio of the α single phase region to the L cross section is 1% or more.
In contrast to this, for example, in the case of the insufficient decarburized region as in a comparative example 201, in the case of using no metal for the second layer as in a comparative example 203, and in the case of not heating to a temperature of the A3 point or higher as in a comparative example 204, it was not possible to obtain a high magnetic flux density in the in-plane circumferential direction thoroughly as in the present invention examples. Further, even when the temperature was increased to a higher temperature and the holding time was made longer as in present invention examples 228 and 229, the similar effect was able to be obtained, but the significant effect did not appear.
Example 4
In this example, as the ferrite-forming element, Sn, Al, Si, Ti, Ga, Ge, Mo, V, Cr, or As was applied to the second layer, and the relationship between the case where demanganization was performed in addition to decarburization and an accumulation degree of {200} planes was examined.
First, base metal sheets containing six types of components F to K shown in Table 9 below and having a balance being composed of Fe and inevitable impurities were prepared. Ingots were each melted by vacuum melting to then be worked to a predetermined thickness by hot rolling and cold rolling, and the above-described base metal sheets were obtained. Incidentally, the A1 point of each of these base metal sheets was 727° C.







TABLE 9







(MASS %)
















STEEL
Ar3 POINT










TYPE
° C.
C
Si
Mn
Al
P
N
S
O



















F
877
0.03
0.05
0.15
0.0005
0.0001
0.0002
<0.0004
0.0002


G
880
0.03
0.10
0.25
0.0004
0.0002
0.0001
<0.0005
0.0001


H
867
0.05
0.05
1.00
0.0003
0.0001
0.0002
<0.0004
0.0002


I
771
0.50
0.1
0.30
0.0004
0.0002
0.0002
<0.0004
0.0002


J
773
0.80
1.00
0.12
0.0030
0.0020
0.0001
<0.003
0.0001


K
859
0.10
0.30
1.50
0.0030
0.0020
0.0001
<0.003
0.0002









In the hot rolling, the ingots each having a thickness of 230 mm were heated to 1000° C. to be thinned down to a thickness of 50 mm, and hot-rolled sheets were obtained. Then, sheet materials having various thicknesses were cut out from these hot-rolled sheets by machining to then be subjected to the cold rolling, and the base metal sheets each having a thickness falling within a range of 10 μm to 750 μm were manufactured.
At this time, the main phase of each of the obtained base metal sheets at room temperature was an α-Fe phase. Further, as a result of measurement, the A3 point at which the α-γ transformation occurred was temperatures shown in Table 9. Further, by X-ray diffraction, a texture in the α-Fe phase of each of the base metal sheets was measured, and by the previously described method, an accumulation degree of {200} planes and an accumulation degree of {222} planes were obtained. As a result, it was confirmed that at the stage of completion of the cold rolling, of each of the base metal sheets, the accumulation degree of the {200} planes was 19 to 27% and the accumulation degree of the {222} planes was 18 to 25%.
Next, these base metal sheets after this cold rolling each had a material promoting decarburization, or a material promoting decarburization and a material promoting demanganization applied thereto as an annealing separating agent, and were subjected to tight coil annealing or stacked annealing. At this time, the annealing was performed so that depths of the decarburization and the demanganization might become not less than 1 μm nor more than 49 μm. As conditions of the annealing, the temperature was set to 700° C. to 900° C. and the annealing was performed in a reduced pressure atmosphere. Further, a structure and a crystal orientation of a surface layer after completion of the decarburization annealing or the decarburization and demanganization annealing were examined. The measurement of the crystal orientation was performed by the X-ray diffraction method, and the accumulation degree of the {200} planes in the α-Fe phase and the accumulation degree of the {222} planes in the α-Fe phase were obtained.
Next, with respect to each of the base metal sheets after the decarburization annealing or the decarburization and demanganization annealing, both surfaces of each of the base metal sheets were coated with the different metal by using an IP method, a hot dipping method, or a sputtering method to have a thickness of 10 μm in total.
Subsequently, a heat treatment was performed under various conditions by the same method as that used in Example 3, and an experiment was performed in which the state in each of the processes during the manufacture was evaluated. An alloyed ratio of the second layer was defined as (S0−S)/S0×100 similarly to Example 3, and assuming that a metal element of the second layer was [M], an area of a region satisfying Fe≦0.5 mass % and [M]≧99.5 mass % was obtained, which was applied to any one of the elements.
On the other hand, a ratio of the α single phase region was also obtained by the same procedure as that in Example 3. However, when the second layer was Sn, T was obtained from an area of a region satisfying Sn≧3.0 mass %, and similarly, in the case of Al, it was obtained from an area of a region satisfying Al≧0.9 mass %. Further, in the case of Si, it was obtained from an area of a region satisfying Si≧1.9 mass %, and in the case of Ti, it was obtained from an area of a region satisfying Ti≧3.0 mass %. Similarly, in the case of Ga, it was obtained from an area of a region satisfying Ga≧4.1 mass %, in the case of Ge, it was obtained from a region satisfying Ge≧6.4 mass %, in the case of Mo, it was obtained from a region satisfying Mo≧3.8 mass %, in the case of V, it was obtained from a region satisfying V≧1.8 mass %, in the case of Cr, it was obtained from a region satisfying Cr≧14.3 mass %, and in the case of As, it was obtained from an area of a region satisfying As≧3.4 mass %.
Table 10 and Table 11 show the base metal sheets and conditions of the heat treatment such as the decarburization, and show the accumulation degrees of the {200} planes and the accumulation degrees of the {222} planes measured during the manufacture (after the decarburization•demanganization annealing) and after the manufacture (after the diffusion treatment), the Z values of the obtained Fe-based metal sheets, the alloyed ratios of the second layers, and evaluation results of the magnetometry.














TABLE 10










DECARBURIZATION









AND
DECARBURIZED



BASE

DEMANGANIZATION
AND
C CONTENT

TEMPERATURE



MATERIAL
SHEET
ANNEALING
DEMANGANIZED
AFTER
FERRITE-
INCREASING



STEEL
THICKNESS
TEMPERATURE
REGION
DECARBURIZATION
FORMING
RATE


No.
TYPE
μm
° C.
μm
mass %
ELEMENT
° C./sec





236
F
150
800
21
0.008
Sn
0.5


237
G
150
800
26
0.010
Sn
0.5


238
H
150
800
23
0.009
Sn
0.5


239
I
150
800
24
0.011
Sn
0.5


240
J
150
800
21
0.009
Sn
0.5


241
K
150
800
26
0.009
Sn
0.5


242
F
10
800
4
0.010
Al
0.5


243
F
100
800
12
0.011
Al
0.5


244
F
250
800
14
0.015
Al
0.5


245
F
500
800
22
0.018
Al
0.5


246
F
750
800
31
0.018
Al
0.5


247
G
150
700
10
0.008
Al
0.5


248
C
150
900
24
0.017
Al
0.5


249
H
200
800
49
0.017
Al
0.5


250
H
200
800
6
0.014
Al
0.5


251
I
100
800
14
0.006
Al
0.5


252
I
100
800
15
0.014
Si
0.5


253
I
100
800
15
0.014
Zn
0.5


254
I
200
800
18
0.008
Ti
0.5


255
I
200
800
18
0.009
Ga
0.5


256
I
200
800
19
0.008
Ge
0.5


257
I
200
800
18
0.010
Ho
0.5


258
I
200
800
17
0.008
V
0.5


259
I
200
800
18
0.009
Cr
0.5


260
I
200
800
18
0.009
As
0.5


261
J
200
800
16
0.010
Al
0.1


262
J
150
800
15
0.007
Al
1




















ACCUMULATION
ACCUMULATION
ACCUMULATION






DEGREE OF
DEGREE OF
DEGREE OF



HOLDING
HOLDING
COOLING
{200} PLANES
{222} PLANES
{200} PLANES



TEMPERATURE
TIME
RATE
AFTER
AFTER
AFTER


No.
T1 ° C.
MINUTE
° C./sec
ANNEALING
ANNEALING}
DIFFUSION





236
1000
5
100
36
21
61


237
1000
5
100
34
25
68


238
1000
5
100
33
21
65


239
1000
5
100
35
23
71


240
1000
5
100
37
21
59


241
1000
5
100
34
18
60


242
1000
5
100
35
21
47


243
1000
5
100
34
24
49


244
1000
5
100
35
18
51


245
1000
5
100
31
26
44


246
1000
5
100
39
22
40


247
1000
5
100
24
26
46


248
1000
5
100
36
18
68


249
1000
5
100
38
27
72


250
1000
5
100
23
20
39


251
1000
5
100
29
16
40


252
1000
5
100
34
24
46


253
1000
5
100
33
26
51


254
1000
5
100
36
22
49


255
1000
5
100
37
24
56


256
1000
5
100
36
23
55


257
1000
5
100
35
26
61


258
1000
5
100
36
19
48


259
1000
5
100
36
28
60


260
1000
5
100
37
15
53


261
1000
5
100
35
22
57


262
1000
5
100
38
17
61




















ACCUMULATION










DEGREE OF




{222} PLANES




AFTER
(S0 − S)/


B50/Bs



No.
DIFFUSION
S0 × 100
T/T0 × 100
Z
OF PRODUCT
ΔB T
NOTE







236
14
66
38
126
0.921
0.021
INVENTION










EXAMPLE 232



237
12
64
33
135
0.942
0.034
INVENTION










EXAMPLE 233



238
13
59
36
129
0.937
0.037
INVENTION










EXAMPLE 234



239
8
61
37
187
0.963
0.007
INVENTION










EXAMPLE 235



240
17
65
34
113
0.921
0.048
INVENTION










EXAMPLE 236



241
12
63
35
123
0.917
0.042
INVENTION










EXAMPLE 237



242
17
79
64
64
0.884
0.069
INVENTION










EXAMPLE 238



243
13
65
50
56
0.879
0.072
INVENTION










EXAMPLE 239



244
11
52
43
63
0.892
0.062
INVENTION










EXAMPLE 240



245
19
39
32
52
0.876
0.075
INVENTION










EXAMPLE 241



246
18
37
30
23
0.857
0.086
INVENTION










EXAMPLE 242



247
14
66
49
35
0.872
0.082
INVENTION










EXAMPLE 243



248
9
64
44
142
0.924
0.027
INVENTION










EXAMPLE 244



249
4
31
26
167
0.947
0.011
INVENTION










EXAMPLE 245



250
15
68
57
5.4
0.853
0.101
INVENTION










EXAMPLE 246



251
17
69
59
28
0.867
0.087
INVENTION










EXAMPLE 247



252
16
66
54
37
0.879
0.083
INVENTION










EXAMPLE 248



253
13
67
56
67
0.896
0.067
INVENTION










EXAMPLE 249



254
19
61
50
29
0.887
0.076
INVENTION










EXAMPLE 250



255
8
59
48
121
0.897
0.031
INVENTION










EXAMPLE 251



256
8
62
51
116
0.916
0.036
INVENTION










EXAMPLE 252



257
6
60
46
129
0.916
0.026
INVENTION










EXAMPLE 253



258
17
58
50
59
0.897
0.068
INVENTION










EXAMPLE 254



259
9
59
48
129
0.906
0.039
INVENTION










EXAMPLE 255



260
11
60
44
119
0.899
0.042
INVENTION










EXAMPLE 256



261
10
57
46
125
0.914
0.029
INVENTION










EXAMPLE 257



262
9
29
24
131
0.916
0.013
INVENTION










EXAMPLE 258
























TABLE 11










DECARBURIZATION









AND
DECARBURIZED



BASE

DEMANGANIZATION
AND
C CONTENT

TEMPERATURE



MATERIAL
SHEET
ANNEALING
DEMANGANIZED
AFTER
FERRITE-
INCREASING



STEEL
THICKNESS
TEMPERATURE
REGION
DECARBURIZATION
FORMING
RATE



TYPE
μm
° C.
μm
mass %
ELEMENT
° C./sec





263
J
150
800
14
0.006
Al
5


264
J
150
800
16
0.007
Al
10


265
J
150
800
14
0.007
Al
20


266
K
150
800
14
0.007
Al
0.5


267
K
150
800
15
0.007
Al
0.5


268
K
150
800
16
0.006
Al
0.5


269
K
300
800
21
0.011
Al
0.5


270
K
300
800
22
0.009
Al
0.5


271
K
300
800
22
0.009
Al
0.5


272
K
300
800
21
0.008
Al
0.5


273
K
300
800
23
0.009
Al
0.5


274
K
300
800
21
0.008
Al
0.5


275
G
300
800
21
0.008
Al
0.5


276
G
300
800
21
0.008
Al
1.5


277
G
300
800
21
0.008
Al
2.5


278
F
8
800
8
0.010
Al
0.5


279
G
100
650
21
0.050
Al
0.5


280
G
100
950
41
0.003
Al
0.5


281
H
100
800
1
0.010
Al
0.5


282
H
200
900
69
0.011
Al
0.5


283
K
100
800
26
0.009
NONE
0.5


284
I
100
800
25
0.009
Al
0.5


285
I
100
800
24
0.010
Al
0.5


286
J
100
800
26
0.009
Al
0.5


287
J
100
800
24
0.008
Al
0.5


288
J
100
800
26
0.010
Al
0.5




















ACCUMULATION
ACCUMULATION
ACCUMULATION






DEGREE OF
DEGREE OF
DEGREE OF



HOLDING
HOLDING
COOLING
{200} PLANES
{222} PLANES
{200} PLANES



TEMPERATURE
TIME
RATE
AFTER
AFTER
AFTER



T1 ° C.
MINUTE
° C./sec
ANNEALING
ANNEALING
DIFFUSION





263
1000
5
100
39
23
75


264
1000
5
100
30
19
42


265
1000
5
100
29
14
38


266
950
5
100
29
24
40


267
1250
5
100
28
23
36


268
1000
0.5
100
30
25
67


269
1000
10
100
22
29
43


270
1000
30
100
21
30
41


271
1000
60
100
22
28
38


272
1000
120
100
28
21
64


273
1000
550
100
27
19
73


274
1000
4500
100
22
31
79


275
1000
5
0.1
29
25
51


276
1000
5
10
26
24
52


277
1000
5
450
24
25
45


278
950
1
100
26
23
47


279
1000
10
100
17
19
21


280
1000
10
100
14
14
19


281
1000
10
100
11
24
23


282
1000
10
100
24
23
57


283
1000
10
100
27
25
12


284
765
10
100
23
29
25


285
1350
10
100
25
16
48


286
1000
6050
100
28
24
46


287
1000
10
0.05
22
26
38


288
1000
10
500
26
22
39




















ACCUMULATION










DEGREE OF




{222} PLANES




AFTER
(S0 − S)/


B50/Bs




DIFFUSION)
S0 × 100
T/T0 × 100
Z
OF PRODUCT
ΔB T
NOTE







263
7
61
47
189 
0.975
0.006
INVENTION










EXAMPLE 259



264
16
76
55
43
0.864
0.064
INVENTION










EXAMPLE 260



265
16
81
68
16
0.846
0.098
INVENTION










EXAMPLE 261



266
11
96
75
21
0.853
0.092
INVENTION










EXAMPLE 262



267
17
100
74
  8.3
0.839
0.103
INVENTION










EXAMPLE 263



268
7
100
76
164 
0.943
0.010
INVENTION










EXAMPLE 264



269
15
79
64
53
0.872
0.059
INVENTION










EXAMPLE 265



270
18
51
43
43
0.867
0.063
INVENTION










EXAMPLE 266



271
16
55
37
12
0.843
0.096
INVENTION










EXAMPLE 267



272
5
56
34
158 
0.929
0.009
INVENTION










EXAMPLE 268



273
3
53
38
168 
0.968
0.007
INVENTION










EXAMPLE 269



274
7
55
42
198 
0.978
0.005
INVENTION










EXAMPLE 270



275
4
54
38
123 
0.895
0.036
INVENTION










EXAMPLE 271



276
6
51
44
128 
0.896
0.034
INVENTION










EXAMPLE 272



277
12
53
31
73
0.879
0.053
INVENTION










EXAMPLE 273



278
19
100
100
86
0.876
0.049
INVENTION










EXAMPLE 274



279
28
65
41
  1.3
0.778
0.123
COMPARATIVE










EXAMPLE 205



280
22
62
38
  0.8
0.779
0.113
COMPARATIVE










EXAMPLE 206



281
14
66
31
  0.9
0.782
0.109
COMPARATIVE










EXAMPLE 207



282
18
61
37
135 
0.905
0.037
INVENTION










EXAMPLE 275



283
11
0
0
  0.6
0.765
0.096
COMPARATIVE










EXAMPLE 208



284
31
57
21
  1.2
0.786
0.109
COMPARATIVE










EXAMPLE 209



285
22
85
70
68
0.875
0.052
INVENTION










EXAMPLE 276



286
15
92
73
63
0.881
0.054
INVENTION










EXAMPLE 277



287
17
84
69
12
0.852
0.089
INVENTION










EXAMPLE 278



288
16
63
31
21
0.859
0.086
INVENTION










EXAMPLE 279










In each of present invention examples, it was possible to confirm that the magnetic flux density difference ΔB becomes a small value as compared to comparative examples and a high magnetic flux density is obtained thoroughly in the in-plane circumferential direction. Further, in these Fe-based metal sheets, it was possible to confirm that an excellent magnetic property in which the value of B50/Bs is 0.80 or more is obtained.
Further, in the present invention examples, as shown in Table 10 and Table 11, it was possible to confirm that the {200} plane in the α-Fe phase is likely to be highly accumulated at each of the stages of the heat treatment.
Further, an L cross section of each of the present invention examples was observed, and thereby it was confirmed that the α single phase region made of the α single phase based component exists in at least a partial region including the surfaces and the ratio of the α single phase region to the L cross section is 1% or more.
In contrast to this, for example, in the case of the insufficient decarburized and demanganized region as in a comparative example 207, in the case of using no metal for the second layer as in a comparative example 208, and in the case of not heating to a temperature of the A3 point or higher as in a comparative example 209, it was not possible to obtain a high magnetic flux density in the in-plane circumferential direction thoroughly as in the present invention examples, and consequently, an obtained magnetic property was also poor. Even when the temperature was increased to a higher temperature and the holding time was made longer as in present invention examples 276 and 277, the similar effect was able to be obtained, but the significant effect did not appear.
In the foregoing, the preferred embodiments of the present invention have been described in detail, but the present invention is not limited to such examples. It is apparent that a person having common knowledge in the technical field to which the present invention belongs is able to devise various variation or modification examples within the range of technical ideas of the present invention, and it should be understood that they also belong to the technical scope of the present invention as a matter of course.
Industrial Applicability
The Fe-based metal sheet of the present invention is suitable for magnetic cores and the like of transformers and the like using a silicon steel sheet, and can contribute to downsizing of these magnetic cores and reduction in energy loss.

## Claims

1. An Fe-based metal sheet, comprising:
less than 0.2 mass % C and having a composition that is capable of causing an α-γ transformation, wherein
a ferrite-forming element, being one element or more selected from the group consisting of Al, Mo, Ga, Sn, Ti, Ge, Sb, V, W, Zn and As is alloyed on a partial or whole region of the Fe-based metal sheet, and
when intensity ratios of respective {001}<470>, {116}<6 12 1>, and {223}<692>directions in a sheet plane by X-ray diffraction are set to A, B, and C respectively and Z=(A +0.97B)/0.98C is satisfied, a Z value is not less than 2.0 nor more than 200;
wherein at least a partial region including surfaces of the Fe-based metal sheet is an α single phase region that is alloyed with said ferrite-forming element, and a ratio of the α single phase region to a cross section of the Fe-based metal sheet is 1% or more,
wherein an accumulation degree of {200} planes is not less than 30% nor more than 99%, and an accumulation degree of {222} planes is not less than 0.01% nor more than 30%, and
said accumulation degree of {200} planes is expressed by Expression (1) below and said accumulation degree of {222} planes is expressed by Expression (2) below,


accumulation degree of {200} planes=[{i(200)/I(200)}/Σ{i(hkl)/I(hkl)}]×100  (1)



accumulation degree of {222} planes=[{i(222)/I(222)}/Σ{i(hkl)/I(hkl)}]×100  (2)


wherein i (hkl) is an actually measured integrated intensity of {hkl} planes in the surface of the Fe-based metal sheet, and I (hkl) is a theoretical intergrated intensity of the {hkl} planes in a sample having a random orientation, and 11 kinds of planes of {110}, {200}, {211}, {310}, {222}, {321}, {411}, {420}, {332}, {521}, and {442} are used as the {hkl} planes.

2. The Fe-based metal sheet according to claim 1, wherein
a layer containing said ferrite-forming element is formed on at least one side of surfaces of the Fe-based metal sheet, and said ferrite-forming element that has diffused from part of the layer is alloyed with Fe.

3. The Fe-based metal sheet according to claim 2, wherein
a thickness of the layer containing said ferrite-forming element is not less than 0.01 μm nor more than 500 μM.

4. The Fe-based metal sheet according to claim 1, wherein a thickness of the Fe-based metal sheet is not less than 10 μm nor more than 6 mm.

5. The Fe-based metal sheet according to claim 1, wherein the α single phase region is formed on a front surface side and a rear surface side of the Fe-based metal sheet, and a crystal grain straddling the a single phase region on the front surface side and the α single phase region on the rear surface side is formed.