JP2014240516A - Nanocrystal soft magnetic alloy and magnetic component using the same - Google Patents

Abstract

PROBLEM TO BE SOLVED: To improve iron loss by selectively improving a conventional composition of a nanocrystal soft magnetic alloy and suppressing a coarse crystal grain phase which becomes a barrier due to pinning against magnetization reversal.SOLUTION: A nanocrystal soft magnetic alloy is composed of a structure including fine crystal grains, represented by compositional formula of FeCuBSiSn, in which x, y, z, and d are given by atom%, and exist in ranges of 0.6≤x≤1.6, 7≤y≤20, 0<z≤16, 0.005≤d≤0.2, and 7≤y+z≤24 and having an average crystal grain size of 60 nm or less, dispersed in an amorphous base phase with a volume fraction of 30% or more. The alloy has: saturation flux density of 1.7T or more, coercive force of 8A/m or less, iron loss of 0.30 W/Kg or less at 1.5T and 50 Hz, and iron loss of 250 W/Kg or less at 1.0T and 10 kHz; a ratio of residual magnetic flux density Bto flux density Bat 80A/m of less than 0.9; and a ratio of the flux density Bat 80A/m to the flux density Bat 800A/m of 0.92 or more.

Description

  The present invention relates to a nanocrystalline soft magnetic alloy having high-saturation magnetic flux density suitable for magnetic core materials such as various reactors, choke coils, pulse power magnetic parts, transformers, motors, and current sensors, and excellent high-frequency magnetic characteristics. It relates to the magnetic parts used.

  Soft magnetic materials used for various reactors, choke coils, pulse power magnetic components, transformers, motor or generator magnetic cores, current sensors, magnetic sensors, antenna cores, electromagnetic wave absorbing sheets, etc. There are crystalline soft magnetic alloys, Fe-based amorphous soft magnetic alloys and Fe-based nanocrystalline soft magnetic alloys. Among these, silicon steel is inexpensive and has a high magnetic flux density, but high frequency has a large loss, and ferrite has a low saturation magnetic flux density, so it is not suitable for high energy density applications with a large operating magnetic flux density. Co-based amorphous soft magnetic alloys are expensive and have a low saturation magnetic flux density of 1 T or less, so when used in high energy density applications, the parts become large and the temperature is unstable due to thermal instability. In applications that are in a high state, there is a problem that loss increases due to aging. The Fe-based amorphous soft magnetic alloy has a low saturation magnetic flux density of about 1.5 to 1.6 T, and it cannot be said that the high-frequency magnetic characteristics are sufficiently superior compared to Co-based amorphous alloys.

Among Fe-based nanocrystalline soft magnetic alloys is one disclosed in Patent Document 1. This alloy has the composition formula: Fe _100-xyz Cu _x B _y X _z (where X is at least one element selected from the group consisting of Si, S, C, P, Al, Ge, Ga, Be) , X, y, and z are atomic%, respectively, and are numbers represented by the following conditions: 0.1 ≦ x ≦ 3, 10 ≦ y ≦ 20, 0 <z ≦ 10, and 10 <y + z ≦ 24)) This is an Fe-based nanocrystalline soft magnetic alloy having a structure in which crystal grains having a grain size of 60 nm or less are dispersed in an amorphous matrix at 30 volume% or more and having a high saturation magnetic flux density of 1.7 T or more.

International Publication WO2007 / 032531

  According to the Fe-based nanocrystalline soft magnetic alloy of Patent Document 1, a high saturation magnetic flux density and a low coercive force can be obtained. When these Fe-based nanocrystalline soft magnetic alloys are mass-produced by an ultra-quenching apparatus such as a single roll method, it is necessary to manufacture a wide alloy ribbon for a long time. At this time, the temperature of the alloy ribbon tends to rise as the production time becomes longer than that of a laboratory-level ribbon production apparatus due to deterioration of the surface properties of the cooling roll or an increase in the roll surface temperature. In addition, the difference in cooling in the width direction of the thin ribbon produced by widening the thin ribbon becomes obvious. For this reason, Cu contained in the alloy is easily diffused and segregated on the surface of the ribbon, and crystals are formed on the outermost surface of the alloy before the heat treatment, which causes surface crystallization and affects the properties. Moreover, the method of Cu segregation differs in the ribbon width direction, which causes variation in characteristics.

  Furthermore, the Cu concentration is decreased in the vicinity of the surface of the Cu segregation part. When heat treatment for nanocrystallization is performed on this, the crystal on the outermost surface grows, a crystal grain layer is formed, and the Cu concentration immediately inside thereof is further reduced. For this reason, the grain size of the crystal grains formed in this region becomes large, preventing magnetization rotation and domain wall movement and deteriorating magnetic saturation, leading to an increase in coercive force and a high-frequency magnetic characteristic. That is, this coarse crystal grain increases the effective magnetic anisotropy in this region, making it difficult for magnetization rotation to occur or pinning the domain wall to make it difficult for the domain wall to move. I guess it has deteriorated. When magnetic saturation deteriorates, there is a problem that the actual magnetic component cannot increase the design magnetic flux density (operating magnetic flux density) and the size of the component increases. Also, if the domain wall is pinned and difficult to move, when the domain wall starts moving from the pinned position, the domain wall moving speed of magnetization reversal increases and eddy current loss increases, resulting in degradation of high frequency magnetic characteristics such as iron loss at high frequencies. To do.

In the Fe-based nanocrystalline soft magnetic alloy disclosed in Patent Document 1, as a means of suppressing coarse crystal grains when manufactured by a mass production machine, a heat treatment that rapidly raises temperature is used to suppress a decrease in the number density of crystal nuclei. Many crystal grains are formed and the crystal grains are refined. However, if the size of the magnetic component is large, rapid temperature increase may not provide a sufficient rate of temperature increase, a temperature difference may occur depending on the position of the component, and uniform heat treatment cannot be performed, and the characteristics may vary depending on the location. There is. In addition, there is a problem that soft magnetism cannot be obtained if the temperature rises too much due to heat generated by the crystal, and it is necessary to mitigate the influence of the temperature rising rate.
Substitution of Nb, Mo 2, etc., which suppresses crystal grain growth, is effective for relaxing the temperature increase rate condition, and Fe—Cu—Nb—Si—B nanocrystalline soft magnetic alloys have been put into practical use. However, since Nb and Mo have a large atomic weight and are non-magnetic, substitution of a large amount of these elements causes a decrease in saturation magnetic flux density. After all, the saturation magnetic flux density of Fe-Cu-Nb-Si-B alloys is 1.6T or less, and most of them are 1.4T or less in practical use, and there is a limit to miniaturization of magnetic parts. Therefore, it is industrially significant that the heating rate condition is relaxed in a high saturation magnetic flux density alloy containing no Nb or Mo.

  Accordingly, an object of the present invention is to suppress the formation of coarse crystal grains and the pinning phenomenon of the domain wall when manufactured by a mass production apparatus, exhibit good magnetic saturation and excellent high-frequency magnetic characteristics, and have a small characteristic variation. A magnetic alloy and a high-performance magnetic part using the same are provided.

The present invention relates to an Fe-Cu-B-Si-based Fe-based nanocrystalline soft magnetic alloy (hereinafter, sometimes referred to as nanocrystalline soft magnetic alloy, nanocrystalline alloy, or simply alloy) in a trace amount and in an appropriate amount range. By adding Sn, the surface crystallization and the formation of coarse grains near the surface are suppressed, the pinning phenomenon of the domain wall is suppressed, the magnetic saturation is excellent, and the high frequency magnetic property is obtained. The inventors have found that a soft magnetic alloy with a small variation due to can be obtained and have arrived at the present invention.
That is, the present invention relates to Fe _100-x-y-zd Cu _x B _y Si _z Sn _d, where x, y, z and d are atomic%, and 0.6 ≦ x ≦ 1.6, 6 ≦ y ≦ 20, 0 <Z ≦ 17, 0.005 ≦ d ≦ 0.2, 7 ≦ y + z ≦ 24, and a structure in which fine crystal grains having an average crystal grain size of 60 nm or less are dispersed in an amorphous matrix by 30% or more by volume fraction It is an alloy.

  This alloy can exhibit excellent magnetic properties with a saturation magnetic flux density of 1.7 T or more, a coercive force of 8 A / m or less, and an iron loss at 1.5 T and 50 Hz of 0.30 W / Kg or less. In addition, excellent high frequency magnetic characteristics with an iron loss of 250 W / Kg or less at 10 kHz and 1 T can be exhibited.

This alloy can have a ratio B _r / B _{80 of} residual magnetic flux density B _r to magnetic flux density B ₈₀ at 80 A / m of less than 0.9.
Further, the ratio B ₈₀ / B ₈₀₀ between the magnetic flux density B _{80 at 80} A / m and the magnetic flux density B _{800 at 800} A / m can be 0.92 or more. The B ₈₀ / B ₈₀₀ is preferably 0.95 or more.

In the alloy of the present invention, Fe can be substituted with 4 atomic% or less of P. By substituting P, the amorphous forming ability is improved. However, since the saturation magnetic flux density is decreased, the substitution amount of P is preferably 4 atomic% or less.
Further, Fe can be substituted with 2 atomic% or less of C. By substituting C, the viscosity of the molten alloy is lowered and the surface properties of the alloy are improved. However, if the amount of C exceeds 2 atomic%, the thermal stability is lowered, which is not preferable.

  The present invention is a magnetic component manufactured using the nanocrystalline soft magnetic alloy described above. This magnetic component can be reduced in size and loss, and a high-performance magnetic component can be realized.

  According to the present invention, even when a Fe-Cu-B-Si nanocrystalline soft magnetic alloy is produced on a mass production machine, Cu surface segregation and coarse grains near the surface are suppressed, and domain wall pinning is performed. The phenomenon can be suppressed. As a result, it is possible to provide a nanocrystalline soft magnetic alloy with good magnetic saturation and high frequency characteristics and small characteristic variations, and a high-performance magnetic component using the same.

It is a figure explaining progress of time and a cooling process by a liquid quenching method. It is a structure | tissue observation (TEM) photograph in the Example which added Sn. (A) shows before heat treatment, (b) shows after heat treatment. It is a figure which shows concentration distribution of each element in the Example which added Sn. It is a structure | tissue observation (TEM) photograph in the comparative example which does not add Sn. (A) shows before heat treatment, (b) shows after heat treatment. It is a figure which shows concentration distribution of each element in the comparative example which does not add Sn. It is a BH curve when Sn is added at 0.1 atomic% and when Sn is not added.

First, the present invention will be described taking a manufacturing process as an example.
The nanocrystalline soft magnetic alloy of the present invention can be manufactured using a liquid quenching method such as a single roll method, but the cooling process is roughly divided into three stages. FIG. 1 shows the passage of time and the cooling process (change in phase state) by the liquid quenching method. In the primary cooling process, the molten metal comes into contact with a roll that rotates at high speed, and is cooled to 10 ⁵ to 10 ⁷ ° C / s very quickly in a short time to be in a supercooled state. Therefore, it is in an amorphous state that is a random atomic arrangement. Thereafter, when the secondary cooling process is started, the cooling rate of the alloy becomes about 10 ³ to 10 ⁵ ° C / s due to the contact between the solid phase and the solid phase. At this time, since Cu is insoluble in any element of Fe-B, if the temperature and time allow Cu to diffuse in the secondary cooling process, Cu is considered to segregate on the surface. After that, in the third cooling process, when the ribbon temperature reaches about 100 to 300 ° C., it is peeled off from the roll, so that the solid phase and the gas phase come into contact with each other, and the cooling rate is greatly reduced. Thus, an alloy ribbon having an amorphous phase as a main phase is produced.

  In the Fe-Cu-B-Si nanocrystalline soft magnetic alloy, the initial fine crystal grains are carried by Cu clusters that are heterogeneous nucleation sites of crystals. During the secondary cooling process, Cu is likely to segregate near the surface of the ribbon. By the way, according to Patent Document 1, it is said that the formation of fine crystal grains can be promoted by substituting various elements including Sn at a ratio of 5 atomic% or less of Fe, and an example including 0.5 atomic% is described in the examples. Has been. However, there is no disclosure about dispersion of initial fine crystal grains and suppression of segregation. Therefore, the present inventors tried to produce an alloy containing Sn in an amount of 0.2 atomic% and an alloy containing 0.5 atomic%. As a result, the alloy containing 0.5 atomic% of Sn was extremely brittle and it was difficult to wind up the ribbon. On the other hand, when a very small amount of Sn of 0.2 atomic% or less was added, the Cu segregation part and the region with a low Cu concentration were eliminated, and the ribbon could be wound up. Furthermore, coarse crystal grains after heat treatment were suppressed, and good magnetic properties were obtained. The effect of this very small amount of Sn addition has not been fully clarified, but is considered as follows.

  Sn is an element having a low melting point and a strong bond with Cu. Therefore, it is considered that Cu diffusion is suppressed and segregation to the surface portion of Cu is suppressed. In addition, it is considered that the addition of a very small amount of Sn suppresses the formation of a region with a low Cu concentration and has an effect of uniformly forming a Cu cluster up to the vicinity of the surface. Therefore, in an alloy containing a very small amount of Sn, coarse crystal grains are suppressed and fine crystal grains having a small grain size are formed in a wide range up to the vicinity of the surface. As a result, the generation of coarse crystal grains is suppressed and the domain wall is less likely to be pinned, the domain wall movement and the magnetization rotation are easily performed, the magnetic saturation is improved, and the coercive force is also reduced. In addition, since the pinning of the domain wall is suppressed, the domain wall moving speed is reduced, eddy current is suppressed, and iron loss at high frequency can be reduced.

  In addition, since the surface segregation of Cu is suppressed by the addition of Sn, it becomes difficult to form a region with a low Cu concentration, and even when heat-treating an amorphous single-phase alloy with a small amount of Cu or an alloy with a small initial microcrystal, By including an appropriate amount of Sn, Cu segregation hardly occurs on the surface, and formation of a region with a low Cu concentration is also suppressed. By adding an appropriate amount of Sn in a small amount, the diffusion behavior of Cu changes, so that segregation of Cu can be suppressed not only during quenching but also during heat treatment, so the effect of Sn can be obtained in a wide Cu content range. Furthermore, by including an appropriate amount of a small amount of Sn, fluctuations in the Cu concentration in the Cu segregation part and in a region with low Cu concentration due to the difference in the cooling rate due to the location of the wide alloy ribbon produced by the mass production device can be reduced. Variations in characteristics due to can be suppressed.

  Next, an alloy composition, a manufacturing method, etc. are demonstrated. Of the terms used in the present specification, “initial fine crystal grains” are crystal nuclei precipitated in an amorphous matrix formed by rapidly cooling a molten alloy, and are several to several tens of nanometers by heat treatment. The term “fine crystal grain” means a crystal grain grown by heat treatment from the initial fine crystal grain or the amorphous single phase. In addition, the volume fraction indicates the proportion of the initial fine crystal grains and the amount of fine crystal grains precipitated in three dimensions and the number density, but the volume fraction was determined by a line segment method from a micrograph as described later. Approximate treatment with two-dimensional values. The number density is obtained by visually counting the number of fine crystal grains per unit area.

Alloy composition of the present invention have the general formula: wherein _{Fe 100-x-y-zd} Cu x B y Si z Sn d, x, y, z, d in atomic%, 0.6 ≦ x ≦ 1.6,6 ≦ y ≦ 20, 0 <z ≦ 17, 0.005 ≦ d ≦ 0.2, 7 ≦ y + z ≦ 24. Of course, the above composition may contain inevitable impurities such as S, O, and N. In order to have a high saturation magnetic flux density (high Bs), it is necessary to have a structure having a fine crystal (nanocrystal) of bcc-Fe, and for that purpose, the Fe content needs to be high. A particularly preferable Fe content is 78 atomic% or more.

  This alloy has a Fe-B-Si system as a basic composition, and contains Cu and Cu, which are nucleation elements insoluble in Fe and Fe, and has good affinity for Cu. The total amount x of Cu is 0.6 ≦ x ≦ 1.6. If Cu is less than 0.6 atomic%, the crystal nucleation effect due to Cu is not sufficient, and coarse crystal grains are formed, which deteriorates magnetic saturation and soft magnetic properties. When the amount of Cu exceeds 1.6 atomic%, embrittlement becomes remarkable at the time of alloy preparation, which is not preferable. The lower limit of the amount of Cu is preferably 0.8 atomic%, more preferably 1.0 atomic%. The upper limit is preferably 1.5 atomic%, more preferably 1.4 atomic%.

  Sn suppresses the segregation of Cu and the formation of regions with low Cu concentration near the surface, and Cu clusters are uniformly formed at a high number density. After the heat treatment, the formation of coarse crystal grains is suppressed, and the average crystal grain size of the surface layer (about 100 nm from the outermost surface) is less than 40 nm. Therefore, there is an effect of suppressing the pinning of the domain wall and improving the high frequency characteristics. In addition, when a wide-width alloy is produced with a mass production device, it works to reduce the difference in Cu segregation depending on the location and to reduce the characteristic variation. However, if the amount of Sn exceeds 0.2 atomic%, the alloy is easily embrittled and manufacturing problems such as difficulty in winding the alloy ribbon occur. Moreover, the said effect by Sn cannot be anticipated that Sn amount is less than 0.005 atomic%. Therefore, the Sn amount d is set to 0.005 ≦ d ≦ 0.2. The lower limit of the Sn amount is preferably 0.01 atomic%, more preferably 0.05 atomic%. The upper limit is preferably 0.15 atomic%, more preferably 0.1 atomic%.

  B (boron) is an element that promotes the formation of an amorphous phase. When the amount of B is less than 6 atomic%, it is difficult to obtain an alloy ribbon having an amorphous phase as a main phase, and the crystal grain size increases. If the amount of B exceeds 20 atomic%, the alloy becomes brittle, which is not preferable. Therefore, the B amount y is 6 ≦ y ≦ 20. The lower limit of the amount of B is preferably 10 atomic%, more preferably 12 atomic%. The upper limit is preferably 18 atomic%, more preferably 16 atomic%.

  By adding Si, the temperature at which the Fe—B compound phase having a large magnetocrystalline anisotropy is precipitated can be increased, and the heat treatment temperature can be increased. For this reason, an alloy containing Si can be heat-treated at a higher temperature. Therefore, since the ratio of the microcrystal grains mainly containing Fe can be increased, the saturation magnetic flux density is increased. Moreover, in the alloy containing Si, alteration or discoloration due to oxidation of the ribbon surface can be suppressed. In particular, when the Si amount z is 1 atomic% or more, an oxide layer mainly composed of Si is formed on the surface of the ribbon, and the oxidation of Fe can be sufficiently suppressed. On the other hand, if the Si content z exceeds 17 atomic%, the saturation magnetic flux density (Bs) is significantly lowered, which is not preferable. The lower limit of the Si amount is preferably 2 atomic%, more preferably 3 atomic%. The upper limit is preferably 10 atomic%, more preferably 8 atomic%.

If necessary, a part of Fe may be substituted with at least one element D selected from Mn, V, Cr, Ti, Zr, Nb, Mo, Hf, Ta, and W. The content of element D is preferably less than 2 atomic%, more preferably 1 atomic%. Most preferably, it is 0.01-0.5 atomic%.
D element stabilizes the amorphous phase remaining after heat treatment, suppresses the growth of fine crystal grains with high Fe content, reduces the average grain size of fine grains, and thus has a saturation magnetic flux density Bs and soft magnetic properties Can be improved. Cr, Ti, Nb, and Ta also have an effect of improving corrosion resistance. However, too much of these elements having a large atomic weight is not preferable because the Fe content per unit weight is lowered and soft magnetic properties are deteriorated.

Further, a part of Fe can be substituted with at least one element selected from Co and Ni. By substituting Co and Ni, the magnitude of the induced magnetic anisotropy can be controlled. Co also has the effect of increasing the saturation magnetic flux density. The preferable Ni content is 0.1 to 5 atomic%, and the particularly preferable Ni content is 0.5 to 1 atomic%. If the Ni content is less than 0.1 atomic%, the effect of improving the handleability (cutting and winding workability) is insufficient, and if it exceeds 5 atomic%, the saturation magnetic flux density Bs and the magnetic flux density B _{80 at 80} A / m are descend. The preferable Co content is 20 atomic% or less. Especially preferably, it is 15 atomic% or less. A particularly high saturation magnetic flux density can be obtained in this range.

  Furthermore, a part of Fe may be substituted with at least one element selected from Re, Y, Zn, As, In, Ag, Sb, platinum group elements, Bi, N, O, and rare earth elements. The total content of these elements is preferably 1.5 atomic percent or less, and more preferably 0.5 atomic percent or less. In order to obtain a particularly high saturation magnetic flux density, the total amount of these elements is preferably 1 atomic% or less, and more preferably 0.5 atomic% or less. In particular, Ag has a synergistic effect of adding a very small amount to promote Cu clustering, and the amount of Cu can be reduced by substituting part of Cu with Ag. For example, it is preferable to replace a trace amount of 0.005 to 0.1 atomic%.

The molten alloy is Fe _100-x-y-zd Cu _x B _y Si _z Sn _d , where x, y, z, d are atomic%, 0.6 ≦ x ≦ 1.6, 6 ≦ y ≦ 20, 0 <z ≦ 17, 0.005 ≦ d ≦ 0.2, 7 ≦ y + z ≦ 24.

  Rapid cooling of the molten alloy can be performed by a single roll method. The molten metal temperature is preferably 50 to 300 ° C. higher than the melting point of the alloy. For example, when manufacturing a ribbon having a thickness of several tens of μm on which initial fine crystal grains are precipitated, a molten metal of about 1250 to 1400 ° C. is placed on the cooling roll from the nozzle. It is preferable to be ejected. The atmosphere in the single roll method is usually produced in the atmosphere when the alloy does not contain an active metal, but may be produced in an inert gas (Ar gas or the like), nitrogen gas or vacuum. In the case of an alloy containing an active metal, it is produced in an inert gas (Ar gas, He gas, etc.), nitrogen gas atmosphere or in vacuum.

Since the peripheral speed of the cooling roll by the single roll method is related to the cooling speed of the alloy ribbon, it is important to control. The roll peripheral speed is preferably 15 to 50 m / s, more preferably 20 to 40 m / s, and most preferably 25 to 35 m / s.
Roll material is pure copper with high thermal conductivity, or copper such as Cu-Be, Cu-Cr, Cu-Zr, Cu-Zr-Cr, Cu-Ni-Si, Cu-Co-Be and Cu-Ni-Be Alloys are suitable. In the case of mass production, or when producing a thick and / or wide ribbon, the roll is preferably water-cooled. Since the water cooling of the roll affects the amorphization and the initial crystal volume fraction generated in the alloy during production, it is necessary to maintain the roll cooling capacity (which may be referred to as the cooling rate) from the beginning to the end of casting.

Since the volume fraction of the initial fine crystal grains is affected by the cooling capacity of the roll (roll material, cooling water channel structure, cooling water amount, etc.), these are controlled so as to obtain an appropriate structure.
Moreover, in the casting of a thin strip using the single roll method, the plate thickness, the cross-sectional shape, the surface undulation, etc. are affected by the paddle shape and the fluctuation of the paddle shape. For controlling the paddle, it is effective to control the distance (= gap) between the nozzle and the roll or adjust the tapping pressure. Gap control is performed by monitoring the distance between the roll and nozzle and constantly applying feedback to adjust the sheet thickness, cross-sectional shape, surface undulation, etc. of the alloy ribbon. If the gap is too wide, it will be difficult to produce a ribbon having a good shape, and the amount of precipitation of the initial fine crystal grains will be different due to the difference in cooling rate caused by variations in the plate thickness. It is effective that the preferred gap is 300 μm or less, more preferably 250 μm or less, particularly preferably 200 μm or less.

By blowing gas from the nozzle between the ribbon and the roll, the ribbon is peeled off from the roll and the ribbon is wound. The stripping temperature of the ribbon can be adjusted by changing the position (peeling position) of the gas blowing nozzle, and is generally 170 to 350 ° C, preferably 200 to 340 ° C, more preferably 250 to 330 ° C.
In the case of mass production, the peeled strip is a wide strip and has a large amount, so it is necessary to wind it directly on a reel. For this reason, when the thin ribbon is remarkably embrittled, the thin strip is cut and winding becomes difficult. Therefore, it is necessary to suppress the embrittlement of the alloy thin strip as much as possible.

  In alloys in which initial microcrystals are included in the production of alloy ribbons, ultrafine initial microcrystal grains having an average grain size of 30 nm or less have a structure dispersed in an amorphous matrix at a ratio of less than 30% by volume. . If the volume fraction of the initial microcrystalline grains in the initial microcrystalline alloy exceeds 30% by volume, the ribbon does not have sufficient toughness and it becomes difficult to handle in the subsequent process, so the volume fraction of the initial microcrystalline grains Is preferably less than 30%.

  Further, in the present invention, the initial microcrystal is not necessarily formed in the rapid cooling state, and can be formed during a part of the heat treatment process. In other words, in the fabricated state, even in the case of an alloy ribbon in which no initial microcrystals exist, the Cu concentration is reduced even during the temperature rising process due to the effect that it is difficult to form a region with a small amount of Cu segregation and average Cu concentration By adding a small amount of Sn, the formation of regions with low number density of Cu clusters as the heterogeneous nucleation sites of crystal grains is suppressed, and the Cu clusters are uniformly formed in the alloy. The number of clusters decreases in the area where the number density decreases, and after heat treatment, uniform fine crystal grains are formed in the entire alloy, and an alloy with excellent magnetic saturation and soft magnetic properties has been realized. Conceivable.

On the other hand, in the case of an alloy that does not contain Sn, Cu that has been uniformly distributed in the liquid phase diffuses in the supercooled liquid state and segregates on the surface, and the Cu concentration decreases on the inner side. For this reason, in the region where the Cu concentration is reduced, the number density of Cu clusters decreases as a heterogeneous nucleation site, and when heat treatment is performed, coarse crystal grains are formed in this region, and magnetic saturation, soft magnetic properties and high frequency Deteriorates magnetic properties.
According to the present invention, as compared with a state containing no Sn, segregation on the surface of Cu is less likely to occur during the heat treatment, and the temperature increase rate condition during the heat treatment can be relaxed. .

The heat treatment of the initial microcrystalline alloy of the present invention is usually performed while maintaining a temperature of 350 ° C. or higher and 500 ° C. or lower. The holding time is preferably 30 seconds to 4 hours. The average rate of temperature increase during the heat treatment is preferably from 0.1 to 200 ° C./min, and more preferably from 0.1 to 100 ° C./min. The average cooling rate during the heat treatment is preferably 0.1 to 200 ° C./min, more preferably 0.1 to 100 ° C./min.
The heat treatment of the alloy containing no initial microcrystals of the present invention is usually performed by rapidly heating to a temperature of 400 ° C. or more and 600 ° C. or less and cooling with or without holding. The holding time is preferably 30 minutes or less, more preferably 5 minutes or less, and particularly preferably 30 seconds or less. The average rate of temperature increase during the heat treatment is preferably 200 ° C./min to 10000 ° C./min, more preferably 500 to 10,000 ° C./min. The average cooling rate during the heat treatment is preferably 1000 to 10,000 ° C./min, and more preferably 2000 to 8000 ° C./min.

  The heat treatment atmosphere can be performed in the air, but an inert gas atmosphere such as Ar gas or a nitrogen gas atmosphere is desirable. The dew point of the atmospheric gas used for the heat treatment is preferably −30 ° C. or less, more preferably −60 ° C. or less from the viewpoint of suppressing surface crystallization. Moreover, heat processing can also be performed in a vacuum.

In addition, induced magnetic anisotropy can be imparted by heat treatment in a magnetic field to control the BH loop shape and improve characteristics. The induced magnetic anisotropy can be imparted by applying a magnetic field having a strength sufficient to saturate the soft magnetic alloy during the temperature rise of the heat treatment, during the holding of the maximum temperature and during at least a part of the cooling. The required minimum magnetic field strength differs depending on whether it is applied in the width direction of the ribbon (in the case of an annular magnetic core, the height direction) or in the longitudinal direction (in the case of an annular magnetic core, the circumferential direction). In general, when a demagnetizing field is applied in a large direction, a large magnetic field is required. When a magnetic field is applied in the circumferential direction (magnetic path direction) of the annular magnetic core, a magnetic field of 400 A / m or more is generally applied. When a magnetic field is applied in the height direction of the annular magnetic core, depending on the shape, a magnetic field larger than the saturation of the magnetic core is applied. In general, the magnetic field applied in the height direction of the annular magnetic core is 40 kA / m or more, and for a normal annular magnetic core, 80 kA / m or more is applied. The applied magnetic field may be any of a DC magnetic field, an AC magnetic field, and a pulse magnetic field. An alloy ribbon or magnetic core having a DC hysteresis loop with a high squareness ratio or a low squareness ratio can be obtained by heat treatment in a magnetic field. In the case of heat treatment without applying a magnetic field, the alloy ribbon and magnetic core show a round type DC hysteresis loop with a medium squareness ratio.
In addition, in this alloy, induced magnetic anisotropy can be imparted also by heat treatment while applying tension.

If necessary, an oxide film such as SiO ₂ , MgO, Al ₂ O ₃ may be formed on the alloy ribbon. By forming the oxide film, the inter-layer resistance of the wound magnetic core and laminated magnetic core formed from this alloy is increased, and the eddy current is reduced, so that the iron loss at high frequency can be further reduced. If necessary, a magnetic core made of a thin ribbon can be impregnated with resin and cured to produce a magnetic core.

The heat-treated alloy has a structure in which microcrystal grains having a body-centered cubic (bcc) structure with an average grain size of 60 nm or less are dispersed in an amorphous matrix with a volume fraction of 30% or more. When the average grain size of the fine crystal grains exceeds 60 nm, the soft magnetic properties are significantly deteriorated. When the volume fraction of the microcrystal grains is less than 30%, the amorphous ratio is too large and the saturation magnetic flux density is low. The average grain size of the crystallites after heat treatment is 40
nm or less is preferable, and 30 nm or less is more preferable. A particularly preferable average crystal grain size is 20 nm or less. Further, the volume fraction of the fine crystal grains after the heat treatment is preferably 40% or more, and more preferably 50% or more. With an average particle size of 60 nm or less and a volume fraction of 30% or more, an alloy having lower magnetostriction and superior soft magnetism than an Fe-based amorphous alloy can be obtained. The Fe-based amorphous alloy ribbon with the same composition has a relatively large magnetostriction, but this nanocrystalline soft magnetic alloy with dispersed fine crystal grains mainly composed of bcc-Fe is about 1/2 of the Fe-based amorphous alloy. It has the following low saturation magnetostriction constant λs and good magnetic saturation.

  In the present invention, the same effect can be obtained not only in the alloy ribbon but also in the alloy powder, and the present invention can also include an alloy powder having the same composition. When producing powder by grinding or atomizing method, as with alloy ribbon, Cu does not segregate on the outermost surface when Sn is not included, and a region with reduced Cu concentration is formed near the inner surface, after heat treatment Since the number of heterogeneous nucleation sites is insufficient and coarse crystal grains are formed, soft magnetic properties are deteriorated. By adding a small amount of Sn even in the powder, a fine structure can be obtained even in the range where the initial fine crystals near the surface are insufficient, and the soft magnetic characteristics can be greatly improved. Variations in nanostructure after heat treatment of the powder can be reduced, yield and work efficiency can be greatly improved.

  Another aspect of the present invention is a high-performance magnetic component using the nanocrystalline soft magnetic alloy. The alloy of the present invention can increase the operating magnetic flux density to be designed. For this reason, it is suitable for high energy density applications where high operating magnetic flux density design is required. Magnetic components include, for example, high current reactors such as anode reactors, choke coils for active filters, choke coils for smoothing, laser power supplies, Suitable for pulse power magnetic parts used in accelerators, various transformers, motors or generator cores. Further, this magnetic part has a lower noise and a preferable result as compared with a magnetic part using an Fe-based amorphous alloy having a large λs. In addition, it can be used for a pulse transformer for communication, a current sensor, a magnetic sensor, an antenna magnetic core, an electromagnetic wave absorbing sheet, a magnetic shield sheet, and the like. Also, a plurality of alloy ribbons can be laminated to form a laminated body, and these laminated bodies can be further laminated to form a laminated structure, and then applied as an iron core for a transformer wound in a step wrap or an overlap.

  The present invention will be described in more detail with reference to the following examples, but the present invention is not limited thereto. The peeling temperature, the average grain size and volume fraction of the fine crystal grains, the element concentration, and the magnetic properties were determined by the following methods.

  The temperature of the alloy ribbon when peeling from the cooling roll by nitrogen gas blown from the nozzle was measured with a radiation thermometer (Apiste Co., model: FSV-7000E) to obtain the peeling temperature.

The average grain size of the microcrystalline grains (same as the initial microcrystalline grains) is the major axis D _{L of} n (30 or more) microcrystalline grains arbitrarily selected from a transmission electron microscope (TEM) photograph of each sample. The minor axis D _S was measured and obtained by averaging according to the formula Σ (D _L + D _S ) / 2n. Also, draw an arbitrary straight line of length Lt on the TEM photograph etc. of each sample, find the total length Lc of the part where each straight line intersects with the fine crystal grains, and the ratio of crystal grains along each straight line L _L = Lc / Lt was calculated. This operation was repeated 5 times, and the volume fraction of fine crystal grains was determined by averaging L _L. Here, the volume fraction V _L = Vc / Vt (Vc is the sum of the volume of the fine crystal grains and Vt is the volume of the sample) is approximated as V _L ≒ Lc ³ / Lt ³ = L _L ³ Treated.
For the number density, a TEM photograph of each sample (manufactured by Hitachi, Ltd.
20,000 times), the number of fine crystal grains of about 3 to 5 nm or more that can be visually confirmed was determined per unit area (μm ² ).

Glow discharge emission analysis (GDOES: Glow Discharge) of the concentration distribution of each element from the surface to the inside of the alloy ribbon
Optical Emission Spectroscopy (manufactured by Horiba, Ltd.) was used for measurement. The element concentration was measured as a concentration distribution in the depth direction by examining the emission intensity of each element when the surface was sputtered. Since the light emission intensity of each element is related to the concentration and the sputtering rate, the vertical axis represents the light emission intensity (corresponding to the element concentration) and the horizontal axis represents the sputtering time (corresponding to the depth).

Obtain a BH curve from a 120mm single plate sample using a direct current magnetization automatic recorder (Metron Giken Co., Ltd.) and obtain a magnetic flux density at 80 A / m.
Magnetic flux density at B ₈₀ and 800 A / m
The magnetic flux density B ₈₀₀₀ (substantially the same as the saturation magnetic flux density Bs) and the residual magnetic flux density Br at B ₈₀₀ and 8000 A / m were measured, and B ₈₀ / B ₈₀₀ and B _r / B ₈₀ were obtained. Incidentally, the reason why B ₈₀₀ is taken here is that the saturation of the B ₈₀₀ region tends to be deteriorated in the alloy according to the present invention. This is because the closer the B ₈₀ / B ₈₀₀ ratio is to 1, the better the saturation of this region becomes.
For iron loss, a 120mm single plate sample was measured with an AC magnetic property evaluation device (manufactured by Toei Kogyo) for iron loss P _{1.5 / 50} (W / Kg) and apparent power S (excitation VA) at 1.5 T, 50 Hz. Went.

Example 1
A nanocrystalline soft magnetic alloy ribbon of Fe _bal Cu _1.3 B ₁₂ Si ₄ Sn _0.1 was prepared as follows. For comparison, a Fe _bal Cu _1.3 B ₁₂ Si ₄ nanocrystalline soft magnetic alloy ribbon without addition of Sn was also produced.
A molten alloy (1300 ° C) having the above composition (atomic%) is a copper alloy cooling roll (width: 168mm, peripheral speed: 27m / s, cooling water inlet temperature: about 60 ° C, outlet temperature: about 70 ° C), which is supercooled in the atmosphere, peeled off from the roll at a strip temperature of 250 ° C, and an initial fine phase whose main phase is an amorphous phase with a width of 25 mm, a thickness of about 22 µm, and a length of about 10,000 m. A ribbon of crystalline alloy was prepared.
At this time, even when the Sn amount was 0.1 atomic%, the winding could be performed and the production was completed. In addition, as a result of measuring the average particle size and volume fraction of the initial fine crystal grains at an arbitrary location, the ratio of the initial fine crystal grains having an average particle size of 30 nm or less in the amorphous matrix to less than 30% by volume in both thin ribbons It was confirmed to have a dispersed structure.
Thereafter, a 120 mm single plate sample taken from each ribbon is put into a heat treatment furnace, heated to 410 ° C. in about 15 minutes, and then subjected to low-temperature low-speed heat treatment for 1 hour, thereby thinning the nanocrystalline soft magnetic alloy. A strip was made.

Both samples were subjected to microstructure observation (TEM) of the roll surface before and after heat treatment. FIG. 2 shows an observation image near the surface of an example in which Sn is added at 0.1 atomic%, and FIG. 4 shows an observation image near the surface in a comparative example without addition of Sn. 2 and FIG. 4A show before heat treatment, and FIG. 4B shows after heat treatment, respectively.
Next, regarding the concentration distribution of each element by GDOES, FIG. 3 shows an example in which 0.1 atomic% of Sn was added, and FIG. 5 shows a comparative example without Sn addition.
The average grain size and volume fraction and number density of the initial fine crystal grains before the heat treatment, the volume fraction of the fine crystal grains after the heat treatment, the soft magnetic characteristics, and the like are shown in Tables 1 and 2 described later.

  It is known as a very important factor for the development of soft magnetic properties of nanocrystalline soft magnetic alloys that the fine grain size is small, but similarly the structure is homogeneous and mainly Fe. It is important that the fine crystal grains are densely packed. When FIG. 2 and FIG. 4 are compared with each other, in FIG. 4A in which Sn is not added, a segregation portion in which crystals are present is observed in the vicinity of the surface, though very little. This segregation part is considered to correspond to the peak of Cu concentration from the concentration distribution of FIG. That is, a Cu segregation portion is generated near the surface, and a region having a low Cu concentration exists inside the Cu segregation portion. This is probably because Cu diffused and moved to the outermost surface to form a Cu segregation part. On the other hand, in FIG. 2A of the present invention, no segregation or the like is found near the surface. According to FIG. 3, it can be seen that the Cu concentration peak disappears, the Cu segregation portion does not exist, and the Cu concentration is substantially constant up to the vicinity of the surface. That is, the fluctuation of Cu concentration near the surface was eliminated by the addition of Sn. Further, comparing FIG. 2B and FIG. 4B, it can be seen that in FIG. 2B to which Sn is added, the crystal grains are smaller and uniformly dispersed to the inside. From the above, the addition of a very small amount of Sn has the effect of eliminating the Cu segregation near the surface and making the average Cu concentration almost constant in the depth direction, which makes the number density of Cu clusters in the depth direction. Thus, it was confirmed that there is a function of making the structure uniform and uniform, and suppressing the formation of coarse crystals.

(Example 2)
Nanocrystalline soft magnetic alloy ribbons with different Sn contents for the compositions shown in Table 1 were produced by the same method and conditions as in Example 1. The thickness of the ribbon was in the range of about 17-30 μm so that the cooling rate was matched as much as possible. However, when the Sn content was 0.5 atomic%, the ribbon was so brittle that it was difficult to wind up. On the other hand, when the Sn content is 0.1 atomic% or less, the ribbon at the stage before winding immediately after pouring can be bent by 180 degrees without breaking until the bending radius reaches 0.5 mm or until it comes into close contact. . From these results, it was determined that the ribbon could be wound up if the Sn content was 0.2 atomic% or less.

Next, as a result of measuring the average particle size and volume fraction of the initial fine crystal grains for each sample, the initial fine crystal grains having an average crystal grain size of 30 nm or less were less than 30% by volume in the amorphous matrix in each thin ribbon. It was confirmed to have a structure dispersed at a rate of. In addition, the volume fraction and number density of the initial fine crystal grains before the heat treatment, and the average crystal grain size and volume fraction in the vicinity of the surface and the parent phase part (depth 5 μm) of the fine crystal grains after the heat treatment were measured.
In addition, the BH curve was obtained from a single plate sample of the nanocrystalline soft magnetic alloy after the heat treatment. The magnetic flux density B ₈₀ and B _800, and the saturation magnetic flux density B ₈₀₀₀ and B _r, the coercive force Hc (A / m) and 1.5T, iron loss _{P 1.5 / 50 (W / Kg} ) at 50 Hz, the apparent power S (VA / Kg) was measured.
The above measurement results are shown in Tables 1 and 2. The saturation magnetic flux density B ₈₀₀₀ is described as Bs and marked with * is an example (the same applies hereinafter).

  From Table 1, it can be seen that the average crystal grain size tends to be larger in the vicinity of the surface than in the internal matrix in any alloy composition after heat treatment. However, comparing the case with Sn and the case without Sn, in the case with Sn, the difference in the grain size between the vicinity of the surface and the parent phase is small, and the average crystal grain size itself is also small. It can also be seen from the structure before heat treatment that the number density of the initial fine crystal grains is high and the crystal grain size is small when Sn is contained. As described above, when an appropriate amount of Sn is added, a fine and dense structure is formed from the vicinity of the surface to the inside of the alloy, and an improvement in magnetic saturation and an improvement in soft magnetic characteristics are realized.

Next, the comparative example (No. 1) without Sn has a relatively high coercive force of 16 A / m, but with Sn, the coercive force is reduced to 7.6 A / m even with a very small amount of 0.05 atomic%. the magnetic flux density B ₈₀ is equal to or greater than 1.64T. Furthermore, the case of adding Sn 0.1 atomic%, the coercive force was reduced to 5.0A / m, B ₈₀ became 1.68T. However, when 0.5 atomic% of Sn is added, the coercive force tends to increase, and as described above, there is a problem in terms of productivity due to low toughness. Further, in the comparative example without Sn, the crystal grains are larger, and the coercive force and iron loss are also affected. Therefore, an excellent soft magnetic property was obtained when the Sn content was appropriate and the coercive force and iron loss were low. That is, when the addition amount of Sn is less than 0.5 atomic%, the ratio B ₈₀ / B ₈₀₀ between the magnetic flux density B _{80 at 80} A / m and the magnetic flux density B _{800 at 800} A / m is 0.92 or more, and 1.7 T The above saturation magnetic flux density is maintained, the coercive force is 8 A / m or less, and the iron loss at 1.5 T and 50 Hz can be 0.3 W / Kg or less. In addition, the apparent power S is approximately 0.5 VA / Kg or less when an appropriate amount of Sn is contained. Even when the amount of Cu and 0.6 to 1.6 atomic%, the coercive force by Sn is added is decreased, the magnetic flux density B ₈₀ was confirmed to tend to rise.

Next, FIG. 6 shows the BH curve with and without Sn. A comparative example (No. 1) without Sn is indicated by a dotted line, and an example (No. 3) in which the Sn amount is 0.1 atomic% is indicated by a solid line.
Looking at this BH curve, it can be seen that the BH curve without Sn forms a pin angle by bulging in the high magnetic flux density region to form a so-called pinning site. This pinning site region has strong anisotropy and poor magnetic saturation. Although it is thought that it appears due to the magnetization process of the tissue, the presence of this region differs in the way of decreasing the magnetic flux density below H = 0 A / m in the demagnetization process. In other words, the demagnetization curve is gentle when Sn is entered, whereas it sharply decreases from the pin stop angle s when Sn is not present. Although it is a little difficult to understand in FIG. 6, the dotted line has a corner and is abruptly falling from the pinning s. This means that the moving speed of the domain wall in the magnetization process is increased by a steep amount. Since the eddy current loss Pe is proportional to the change rate of the magnetic flux density B (Pe∝dB / dt), the eddy current loss increases as the change rate of the magnetic flux density dB / dt increases, which results in an increase in iron loss. Leads to. Actually, the iron loss P _{1.5 / 50} of 1.5T, 50Hz is 0.29 W / kg in the example (No. 2) in which 0.05 atomic% of Sn is added, and 0.21 W / kg in the example (No. 3). In the comparative example (No. 1) without Sn, it increased to 0.58 W / kg, and in the comparative example (No. 7), it increased to 0.52 W / kg. The comparative example (No. 5) is comparatively small at 0.31 W / Kg, but this may have a relatively large plate thickness and a high number density of initial microcrystals. In addition, it is reflected in the squareness on the BH curve, and in the comparative example, B ₈₀ / B ₈₀₀ is 0.90 or higher and the saturation is high, but B _r / B ₈₀ is also 0.9 or more and the increase in squareness cannot be suppressed. It is the result. That is, it can be said that the pin stop angle stands and the pin stop action works to prevent the domain wall from moving.

  Further, for the alloys of Comparative Example (No. 1) and Example (No. 2), a ring sample having an outer diameter of 25 mm and an inner diameter of 20 mm was prepared by photoetching of a ribbon and subjected to the same heat treatment. Ten samples were stacked and the high-frequency iron loss at 1 T and 10 kHz was measured. As a result, the iron loss of the example (No. 2) was 98 W / Kg, but the iron loss of the comparative example (No. 1) without Sn was 270 W / Kg. As described above, it was confirmed that the high-frequency magnetic characteristics were greatly improved in the examples.

Example 3
In the same manner as in Example 1, Fe _bal. Cu _x B ₁₂ Si ₄ Sn _d
An amorphous alloy ribbon having a composition (0.6 ≦ x ≦ 1.0, 0 ≦ d ≦ 0.1) was produced. The alloy of this composition was an amorphous single phase. However, although it is amorphous, nuclei of initial fine crystal grains are present. Next, the alloy ribbon was cut to prepare a sample having a width of 25 mm and a length of 120 mm, and subjected to a rapid temperature raising heat treatment at 50 ° C./s in an infrared intensive heating furnace in an argon gas atmosphere. Rapid heat-up heat treatment has an average rate of temperature rise from 300 ° C to the holding temperature (gradient of temperature rise time) of 50 ° C / s
The sample was heat-treated at 450 ° C. for 10 seconds, and then cooled to obtain a heat-treated sample. As a result of measuring the average grain size and volume fraction of the fine crystal grains for each sample, each sample has a structure in which fine crystal grains having an average grain diameter of 60 nm or less are dispersed in an amorphous matrix at a ratio of 30% by volume or more. It was confirmed to have The average particle diameter of the surface layer (about 100 nm from the outermost surface) and the magnetic flux density B ₈₀ , B ₈₀₀ , B ₈₀₀₀ , coercive force Hc and 1.5 T, iron loss P _{1.5 / 50} at ₅₀ Hz were measured.
The above measurement results are shown in Table 3.

As shown in Table 3, in the case of heat treatment at a heating rate of 50 ° C./s, the magnetic properties of the magnetic flux density B ₈₀ , coercive force Hc and B ₈₀ / B ₈₀₀ were insufficient in the comparative example without Sn. In these comparative examples, the number density of the nuclei of the initial fine crystal grains is reduced, few crystal grains are actively grown, and each crystal grain is coarsened. This is presumably because Cu, which was uniformly distributed in the liquid phase, was in a supercooled liquid state and was quenched into an amorphous phase (solid phase) in a state where it began to aggregate. In the rapid cooling preparation state, there is a potential fluctuation of Cu concentration, and it is thought that in the heat treatment process, initial fine crystal grains are generated using them as nuclei. However, in this state in which the supersaturation is not reached, the nucleus is insufficient, the coercive force _Hc is increased, and the intended soft magnetic property cannot be obtained. On the other hand, when Sn was added, high number density nanocrystal grains appeared in a region where the initial fine crystal grains were not present. The addition of Sn suppresses Cu concentration fluctuations in the quenching state and causes Cu diffusion, clustering, nucleation, initial microcrystal precipitation, and grain growth, resulting in high number density nuclei. It is thought that it was done. If the number density of bccFe grains is high, the residual amorphous phase
Fe concentration decreases and the amorphous phase stabilizes, so crystal grain growth is suppressed. The soft magnetic properties are No. 12 than No. 11, and No. 15 than No. 14, Sn and Cu, respectively.
B ₈₀ , Hc, and B ₈₀ / B ₈₀₀ are all improved.

Example 4
Next, the effect of changing the heating rate was examined.
An alloy ribbon having a composition of Fe _bal. Cu _1.0 B ₁₂ Si ₄ and Fe _bal. Cu _1.0 B ₁₂ Si ₄ Sn _0.1 was produced under the same conditions as in Example 1. It was confirmed that each thin ribbon had a structure in which initial fine crystal grains having an average crystal grain size of 30 nm or less were dispersed in an amorphous matrix at a ratio of less than 30% by volume. 10 ℃ / s for this alloy ribbon, 50
A heat treatment was performed at a rate of temperature increase of 450 ° C./s at a rate of 100 ° C./s. The average particle diameter of these surface layers, magnetic flux density B ₈₀ , B ₈₀₀ , B ₈₀₀₀ , coercive force Hc and iron loss P _{1.5 / 50} at _1.5 T, 50 Hz were measured. The results are shown in Table 4.

As shown in Table 4, in the comparative example without Sn, 100 ° C / s
When the heat treatment is performed, the magnetic flux density and the coercive force are improved. On the other hand, in the example containing Sn, the effect of reducing the coercive force is observed even at 10 ° C./s, and the effect of reducing the coercive force is extremely high at 50 ° C./s or more. The coercive force largely depends on the refinement of the structure, but it is derived from the increase in the number density of nuclei by adding Sn. When the diffusion time of Cu until the temperature rise rate is slow and nucleation is long, the number density once reaches a peak and then starts to decrease, so if the temperature rise rate is too slow, the number of nuclei decreases too much. As a result, a fine and high-density nanocrystal grain structure cannot be obtained. However, by containing a small amount of Sn in an appropriate amount, Cu in a rapidly cooled state is prepared.
This makes the distribution of Cu more uniform, so that even when the heating rate is slow, a sufficient number density of Cu clusters is ensured, contributing to the refinement of the structure. Therefore, it has been found that the temperature rise rate dependency is reduced, the temperature rise rate condition during the heat treatment can be greatly improved, and the difficulty of the heat treatment is eliminated.

(Example 5)
An alloy ribbon having a composition shown in Table 5 and a thickness of about 20 to 22 μm and a width of 50 mm was produced in the same manner as in Example 1. Next, sample Nos. 6-1 to 6-14 were produced by slitting the alloy ribbon to a width of 5 mm. At this time, Samples C5 and C25 were prepared so that the center of the slit ribbon was located at a position 5 mm from the end of the ribbon 50 mm in width and a position 25 mm, respectively. Further, Samples C5 and C25 were taken from a position about 100 m from the tip of the ribbon where casting started and a position about 7500 m, respectively. For comparison, a Fe _bal Cu _1.0 Si ₃ B ₁₂ alloy containing no Sn was prepared in the same manner as a comparative example. In addition, the sample slit so that the center is at the 5 mm position is denoted as C5, and the material slit so that the center is also at the 25 mm position is denoted as C25.
For these samples (before heat treatment), the presence or absence of a Cu segregation part was confirmed by glow discharge emission analysis. As a result, Cu segregation occurred in both the C5 and C25 samples of No. 5-15 to 5-18 without Sn. C25 was more prominent. On the other hand, in Nos. 5-1 to 5-14 containing Sn, a remarkable Cu segregation part was not recognized together with the C5 and C25 samples.

Next, these alloy ribbon samples were wound around an outer diameter of 15.5 mm and an inner diameter of 15 mm to produce a wound core, and the same heat treatment as in Example 3 was performed. After the heat treatment, it was confirmed that the crystal grains mainly containing uniform fine bcc Fe having an average crystal grain size of about 15 nm were dispersed in an amorphous matrix at 30 volume% or more. In each sample No.5-1 to 5-18, two wound cores were prepared, one wound core sample analyzed the microstructure and element concentration distribution of the alloy, and the other wound core sample was Saturation magnetic flux density Bs, iron loss P _{1 / 10k} at high frequency of 10 kHz, B ₈₀ / B ₈₀₀ and 1T, which are indicators of magnetic saturation, were measured. The results obtained are shown in Table 5.

  From the results of Table 5, it can be seen that the present invention example containing an appropriate amount of a small amount of Sn has better magnetic saturation than the comparative example containing no Sn and low iron loss in the high frequency region. In addition, it was confirmed that the difference in characteristics depending on the location of the ribbon is small and the characteristic variation is small.

Further, for comparison, the iron loss P _{1 / 10k} of a 6.5 mass% silicon steel with a thickness of 100 μm at 1 T and 10 kHz was measured. P _{1 / 10k} is 600 W / kg. The alloy of the present invention showed a lower high-frequency iron loss value and was confirmed to be excellent in high-frequency characteristics.

(Example 6)
An alloy ribbon having a composition shown in Table 6 and a thickness of 20 to 22 μm and a width of 50 mm was produced in the same manner as in Example 1. Next, the surface roughness of these alloy ribbons was measured. Further, a ring sample having an outer diameter of 25 mm and an inner diameter of 20 mm was produced from these alloy ribbons, and the same heat treatment as in Example 3 was performed. After the heat treatment, it was confirmed that the crystal grains mainly containing uniform fine bcc Fe having an average crystal grain size of about 15 nm were dispersed in an amorphous matrix at 30 volume% or more. Thereafter, the saturation magnetic flux density Bs, B ₈₀ / B ₈₀₀ and 1 T, which are indicators of magnetic saturation, and iron loss P _{1/10 k} at a high frequency of 10 kHz were measured. The results obtained are shown in Table 6.

  From the results of Table 6, the present invention example containing an appropriate amount of a small amount of Sn is excellent in magnetic saturation and low in iron loss at high frequencies. Further, the alloy containing C has a small surface roughness and an improved surface condition. On the other hand, it was confirmed that the comparative example not containing Sn was inferior in magnetic saturation and high-frequency iron loss and inferior in characteristics to the present invention.

(Example 7)
Nanocrystalline soft magnetic alloy ribbons having a constant Sn content and varying Cu, B, Si, etc. for the compositions shown in Table 7 were produced by the same method and conditions as in Example 1. The thickness of the ribbon was in the range of about 17-30 μm so that the cooling rate was matched as much as possible.
The alloy ribbon was subjected to a heat treatment in which the temperature was rapidly raised to 450 ° C. at a rate of 50 ° C./s and held for 1 minute. Each sample was confirmed to have a structure in which fine crystal grains having an average crystal grain size of 60 nm or less were dispersed in an amorphous matrix at a ratio of 30% by volume or more. The average particle diameter of the surface layer of these samples, the magnetic flux densities B ₈₀ and B ₈₀₀₀ , the coercive force Hc, B ₈₀ / B ₈₀₀ and 1.5 T, and the iron loss P _{1.5 / 50} at ₅₀ Hz were measured. The results are shown in Table 7.
From the results in Table 7, it was found that the effect of Sn addition was obtained even when Cu, B, Si, etc. were changed.

In addition, the present invention is intended to express an effective microcrystalline structure using Cu clusters that act as heterogeneous nucleation sites in the Fe-B-Si amorphous matrix, By adding an appropriate amount of Sn, the surface segregation of Cu and the region with low Cu concentration are reduced, and when Cu clusters are uniformly distributed in the alloy and crystallized by heat treatment, nanocrystal grains are uniformly and finely formed. Dispersed in the amorphous matrix to achieve excellent characteristics. The present invention can be applied to any alloy that exhibits the same effect by adding a small amount of Sn.

Claims (9)

Composition formula: Fe _100-x-y-zd Cu _x B _y Si _z Sn _d where x, y, z and d are atomic%, 0.6 ≦ x ≦ 1.6, 6 ≦ y ≦ 20, 0 <z ≦ 17, 0.005 ≦ d ≦ 0.2, 7 ≦ y + z ≦ 24, an alloy having a structure in which fine crystal grains having an average crystal grain size of 60 nm or less are dispersed in an amorphous matrix by 30% or more by volume fraction A nanocrystalline soft magnetic alloy characterized by that. 2. The nanocrystalline soft magnetic alloy according to claim 1, wherein the saturation magnetic flux density is 1.7 T or more, the coercive force is 8 A / m or less, and the iron loss at 1.5 T and 50 Hz is 0.30 W / Kg or less. The nanocrystalline soft magnetic alloy according to claim 1 or 2, wherein an iron loss at 1.0 T and 10 kHz is 250 W / Kg or less. 4. The nanocrystalline soft magnetism according to claim 1, wherein the ratio B _r / B _{80 of the} residual magnetic flux density B _r to the magnetic flux density B ₈₀ at 80 A / m is less than 0.9. 5. alloy. 5. The ratio B ₈₀ / B ₈₀₀ between the magnetic flux density B _{80 at 80} A / m and the magnetic flux density B _{800 at 800} A / m is 0.92 or more, according to claim 1. Nanocrystalline soft magnetic alloy. The nanocrystalline soft magnetic alloy according to any one of claims 1 to 5, wherein Fe is substituted with 4 atomic% or less of P. The nanocrystalline soft magnetic alloy according to any one of claims 1 to 6, wherein Fe is substituted with 2 atomic% or less of C. It is obtained by heat-treating an initial microcrystalline alloy having a structure in which initial microcrystalline grains having an average crystal grain size of 30 nm or less are dispersed in an amorphous matrix at a volume fraction of less than 30%. The nanocrystalline soft magnetic alloy according to any one of claims 1 to 7. A magnetic component manufactured using the nanocrystalline soft magnetic alloy according to claim 1.