KR100366827B1 - high saturated magnetic flux density low iron loss Fe based soft magnetic alloy and magnetic core using same, and method of producing same - Google Patents

Abstract

The present invention is a general formula containing Fe _b Zr _x Nb _y B _β or Zn, wherein in the composition of (Fe _1-a Q _a ) _b B _β M _λ Zn _z , Q is either Co or Ni , Or both, M is one or two or more elements selected from Ti, Zr, Hf, V, Nb, Ta, Mo, W, a≤0.05, b≥80 atomic%, 5 atomic% ≤x + y≤7.5 Satisfying conditions such as atomic%, 1.5 / 6 ≤ x / (x + y) ≤ 2.5 / 6, 5 atomic% ≤ β ≤ 12.5 atomic%, and a highly saturated magnetic flux consisting of microcrystalline grains and the balance of an amorphous alloy phase of bccFe The present invention relates to a density low iron loss Fe base soft magnetic alloy, a low iron loss magnetic core using the same, and a method of manufacturing the same. The alloy of the present invention is excellent in soft magnetic properties such as low coercive force, high permeability, and high saturation magnetic flux density.

Description

High saturated magnetic flux density low iron loss Fe based soft magnetic alloy and magnetic core using same, and method of producing same}

The present invention relates to a high saturation magnetic flux density low iron loss Fe-based soft magnetic alloy used in magnetic heads, trans, choke coils and magnetic cores using the same, and to a method of manufacturing the same, in particular, high saturation magnetic flux density, low iron loss Density low iron loss Fe-based soft magnetic alloy, magnetic core and a method for producing the same.

The characteristics generally required for soft magnetic alloys used for magnetic heads, transformers, and choke coils include high saturation magnetic flux densities, high permeability, low coercive force, and the like. Moreover, especially when used for a trans | transformer, low iron loss is calculated | required.

Therefore, in the case of producing a soft magnetic alloy, research on materials has been made in various alloy systems from these viewpoints.

Conventionally, Fe, Si alloy, FeNi alloy, etc. are used for the above-mentioned use, and the amorphous alloy of Fe group or Co group was used recently.

Although the FeSi alloy has a high saturation magnetic flux density, the iron loss is large, and therefore, there is a problem that the power loss increases when used in a transformer.

In addition, the FeNi alloy has a problem that the saturation magnetic flux density is lowered in an alloy composition having excellent soft magnetic properties.

In addition, the Fe-based amorphous alloy is excellent in saturation magnetic flux density and iron loss, but has low thermal stability, and therefore has a problem that the change in magnetic properties over time is large. In addition, since the Co-based amorphous alloy has a low saturation magnetic flux density, there is a problem in that it is not suitable for a transformer for power conversion.

However, an important characteristic of the soft magnetic alloy for trans is that the iron loss is small and the saturation magnetic flux density is high.

Conventionally, the iron loss of silicon steel widely used as a transformer is 1.0 W / kg (for 1.7 T, 50 kPa) and the saturation magnetic flux density is 2.0 T, and therefore the iron loss must be made smaller.

In addition, the iron loss of the Fe-based amorphous alloy for transformers, which is conventionally used as a part of the use, is 0.25 W / kg (for 1.4 T, 50 kPa), the saturation magnetic flux density is 1.56 T, and furthermore, the iron loss is made smaller and saturated. There is a problem to increase the magnetic flux density.

On the basis of such a background, the present inventors have already filed a patent for the Fe-based soft magnetic alloy of high saturation magnetic flux density in Japanese Patent Laid-Open No. 7-65145, Japanese Patent Laid-Open No. 5-93249, and the like.

One of the alloys related to this patent application was a highly saturated magnetic flux density soft magnetic alloy comprising a composition represented by the following formula.

(Fe _1-a1 Q _a1 ) bBx ₁ Ty ₁

Provided that Q is either or both of Co and Ni, and T is one or two or more elements selected from the group consisting of Ti, Zr, Hf, V, Nb, Ta, Mo, and W; Either or both of Hf contain a ₁ ≦ 0.05, b ≦ 93 atomic%, x ₁ = 0.5 to 8 atomic%, y ₁ = 4 to 9 atomic%.

The other one of the alloys related to the above-mentioned patent application was a high saturation magnetic flux density alloy comprising a composition represented by the following formula.

FebBx ₂ Ty ₂

Provided that T is one or two or more elements selected from the group consisting of Ti, Zr, Hf, V, Nb, Ta, Mo, and W, and contains either or both of Zr and Hf; 93 atomic%, x ₂ = 0.5 to 8 atomic%, y ₂ = 4 to 9 atomic%.

By the way, when it is going to obtain the soft magnetic alloy which has these characteristics, the material which exhibits the target soft magnetic characteristics can be obtained by adding an element newly, or adjusting the composition ratio.

However, in order to determine the properties of the material, it is generally insufficient to make the regulations concerning these compositions, and it is necessary to pay sufficient attention to the manufacturing method. For example, as a post-process in the case of obtaining an alloy in which most of the amorphous phase is obtained by a single-roll liquid quenching method using a crucible equipped with a cooling roll and a nozzle, how to heat-treat the alloy, for example, annealing treatment. By a large influence on the characteristic of the material to be produced. More specifically, it depends on how many times the annealing is actually performed (determination of the heat treatment temperature), and the temperature-time relationship until reaching the heat treatment temperature, that is, what temperature rise rate is used. The properties of the alloy are subject to change.

In the case of obtaining an alloy by a single-roll liquid quenching method using a crucible having a cooling roll and a nozzle, the temperature of the melt when the melt is injected from the nozzle, that is, the injection temperature is produced. It will greatly affect the properties of the material.

In the case where the soft magnetic alloy is used as a transformer or the like, the soft magnetic alloy must be left in a heated state for a long time during processing at the time of manufacture. The soft magnetic alloy prepared as described above is left for a long time in a heated state (high temperature state). There was a problem that the change in magnetic properties over time was large. In addition, when the soft magnetic alloy is used in a small electronic device, for example, a magnetic head, when the normal detection current is raised to further improve the output, the magnetic head becomes a high temperature exceeding about 200 ° C. When used for a long time in a high temperature state, there is a problem in the reliability of the obtained product because the change of the magnetic properties of the soft magnetic alloy with time is large.

As described above, the present inventors have developed various Fe-based soft magnetic alloys of the above-mentioned compositions. As a result of studying the alloys of the above-mentioned compositions, the present inventors found that the composition ratio of Zr and Nb among the elements T in the alloy was determined. When limited to the range, it was found to exhibit excellent soft magnetic properties.

Moreover, although it was not anticipated at all conventionally, by limiting the composition ratio of Zr and Nb, it became clear that iron loss fell largely at several hundred kPa or less, and it became clear also that breakage deformation was very large.

In this way, the present inventors have reached the present invention.

Accordingly, a first object of the present invention is to provide a highly saturated magnetic flux density low iron loss Fe-based soft magnetic alloy or magnetic core having excellent soft magnetic properties, small iron loss, large breakage strain, and excellent workability.

Further, the second object of the present invention is that the iron loss can be 0.10 w / kg (at 1.4 T, 50 kPa) or less, and the saturation magnetic flux density can be 1.5 T or more, so that the magnetic properties can be maintained for a long time in a heated state. It is possible to provide a highly saturated magnetic flux density low iron loss Fe-based soft magnetic alloy or magnetic core, which is less changeable over time and can cope with bending during manufacturing trans, etc.

The third object of the present invention has been made in view of the above circumstances, and an object thereof is to provide a method for producing an alloy excellent in low coercive force, high permeability, and high saturation magnetic flux density characteristics.

The fourth object of the present invention is to develop a method for producing a soft magnetic alloy according to the above patent application, to improve soft magnetic properties, and to change the magnetic properties over time even if left for a long time at a high temperature, It is to provide a method for producing a Fe-based magnetic magnetic alloy, which is preferable for a transformer or the like that can cope with processing during production.

In order to achieve the above object, the present invention adopts the following configuration.

In the first highly saturated magnetic flux density low iron loss Fe-based soft magnetic alloy of the present invention, an alloy mainly composed of an amorphous phase is heat-treated, and at least 50% of the structure is mainly composed of a bcc-Fe phase having an average grain size of 100 nm or less. It becomes a crystalline structure and consists of a composition shown by following formula.

Fe _a Zr _x Nb _y B _z

Where a, x, y, and z are 80 atomic% ≤ a, 5 atomic% ≤ x + y ≤ 7 atomic%, 1.5 / 6 ≤ x / (x + y) ≤ 2.5 / 6, 5 atomic% ≤ z ≤ 12.5 Is atomic percent.

In addition, a, x, y, z representing the above composition ratio is 83 atomic% ≤ a, 5.7 atomic% ≤ x + y ≤ 6.5 atomic%, 1.5 / 6 ≤ x / (x + y) ≤ 2.5 / 6, 6 atomic% ≤ z≤9.5 atomic% may be sufficient.

In addition, a, x, y, z representing the above composition ratios are 85 atomic% ≤ a ≤ 86 atomic%, 5.7 atomic% ≤ x + y ≤ 6.5 atomic%, x / (x + y) = 2/6, 8 atomic% ≤ z≤9 atomic% may be sufficient.

Moreover, it is preferable that x and y which show the composition ratio of Zr and Nb are 1.5 atomic% <= <= 2.5 atomic%, 3.5 atomic% <= y <5.0 atomic%.

The first highly saturated magnetic flux density low iron loss Fe-based soft magnetic alloy of the present invention has the crystallization temperature of the bcc-Fe phase as T _x1 and the crystallization temperature of the compound phase crystallized at a higher temperature than T _x1 as T _x2 . When the interval ΔT _x of the crystallization temperature is ΔT _x = T _x2 -T _x1 , the temperature is 200 ° C. ≦ ΔT _x .

The first highly saturated magnetic flux density low iron loss Fe-based soft magnetic alloy of the present invention has a saturation magnetic flux density of 1.5 T or more.

Further, the first highly saturated magnetic flux density low iron loss Fe-based soft magnetic alloy of the present invention has an iron loss of 0.15 W / kg or less when a magnetic flux of 1.4 T is applied at a frequency of 50 Hz.

In addition, the iron loss change rate of the first highly saturated magnetic flux density low iron loss Fe-based soft magnetic alloy of the present invention before and after aging at 200 ° C. for 500 hours is 10% or less.

In the first highly saturated magnetic flux density low iron loss Fe-based soft magnetic alloy of the present invention, the fracture strain is 1.0 × 10 ⁻² or more.

Moreover, the conditions for heat-treating an alloy are a temperature increase rate of 10 degrees C / min or more, More preferably, 10 degrees C / min or more and 200 degrees C / min or less, More preferably, 30 degrees C / min or more and 100 degrees C / min or less, and heat treatment temperature 490. It is preferable to set it as C or more and 670 degrees C or less, More preferably, it is 500 or more and 560 degrees C or less, and there is no holding time or 1 hour or less.

The first highly saturated magnetic flux density low iron loss Fe-based soft magnetic alloy of the present invention contains Fe, Zr, Nb, and B. In particular, the composition ratio of Zr (x) and Nb (y) is 5.7 atomic% ≦ x + y. By setting it as ≦ 6.5 atomic% and in the range of 1.5 / 6 ≦ x / (x + y) ≦ 2.5 / 6, the magnetic permeability and the saturation magnetic flux density are improved, and the iron loss is reduced. In particular, the reduction of iron loss in the low frequency region of several hundreds of kHz or less is remarkable, and the change over time of iron loss is also small.

In addition, since the composition ratio of Fe, which is an element responsible for magnetic properties, is relatively high at 80 atomic% or more, preferably 83 atomic% or more, more preferably 85 atomic% or more, the saturation magnetic flux density is increased to make the high saturation magnetic flux density low iron loss. It is possible to improve the Fe-based soft magnetic alloy.

In addition, since the composition ratio of B having an amorphous forming ability is 12.5 atomic% or less, preferably 9.5 atomic% or less, more preferably 9 atomic% or less, the FeB-based compound is used when the alloy is heat-treated to precipitate the microcrystalline structure. Is suppressed, and a decrease in soft magnetic properties is prevented.

If the interval ΔT _x of the crystallization temperature is 200 ° C. or more, the interval between the crystallization temperature of the bcc-Fe phase and the compound phase crystallized on the high temperature side of the bcc-Fe phase increases, so that the alloy is heat treated under optimum conditions. This makes it possible to precipitate only the bcc-Fe phase to suppress the precipitation of other compound phases and to improve the soft magnetic properties of the highly saturated magnetic flux low iron loss Fe-based soft magnetic alloy.

In order to solve the above problem, the second highly saturated magnetic flux density low iron loss Fe-based soft magnetic alloy of the present invention is composed of a composition represented by the following formula, and at least 50% or more of the microstructure of bccFe having an average grain size of 100 nm or less It consists of crystal grains, the balance is made of an amorphous alloy phase, the microcrystalline grains of bccFe is to quench the alloy to form an approximately single-phase structure of amorphous phase, and then cooled and precipitated after heating the amorphous phase above the crystallization temperature.

(Fe _1-a Q _a ) _b B _x M _y Zn _z

Provided that Q is either or both of Co and Ni, and M is one or two or more elements selected from Ti, Zr, Hf, V, Nb, Ta, Mo, and W; x, y, and z are 0 ≦ a ≦ 0.05, 80 atomic% ≦ b, 5 atomic% ≦ x ≦ 12.5 atomic%, 5 atomic% ≦ y ≦ 7 atomic% and 0.025 atomic% ≦ z ≦ 0.2 atomic%.

In order to solve the above problem, the second highly saturated magnetic flux density low iron loss Fe-based soft magnetic alloy of the present invention is composed of a composition represented by the following formula, and at least 50% or more of the structure has a bccFe of 100 nm or less in average grain size. Consisting of fine crystal grains of which the remainder consists of an amorphous alloy phase, wherein the bccFe microcrystalline grains are quenched to form an amorphous single phase structure, and then the amorphous phase is cooled and precipitated after heating above the crystallization temperature. .

(Fe _1-a Q _a ) _b B _x M _y Zn _z M ' _u

Provided that M is one or two or more elements selected from Ti, Zr, Hf, V, Nb, Ta, Mo, and W, and M 'is one or two or more elements selected from Cr, Ru, Rh, and Ir. , B, x, y, z, u indicating composition ratio are 0 ≦ a ≦ 0.05, 80 atomic% ≦ b, 5 atomic% ≦ x ≦ 12.5 atomic%, 5 atomic% ≦ y ≦ 7 atomic%, 0.025 atomic% ≤ z ≤ 0.2 atomic%, u ≤ 5 atomic%.

In order to solve the above problem, the second highly saturated magnetic flux density low iron loss Fe-based soft magnetic alloy of the present invention is the iron loss by heating at 320 ° C. for 100 hours in the highly saturated magnetic flux density low iron loss Fe-based magnetic alloy described above. The change rate is 20% or less, the saturation magnetic flux density is 1.5 T or more, and the magnetic permeability is 30,000 or more.

In the second highly saturated magnetic flux density low iron loss Fe-based soft magnetic alloy of the present invention, in order to solve the above problems, the fracture strain is 10 × 10 ⁻³ or more.

The third high saturation magnetic flux density low iron loss Fe-based soft magnetic alloy of the present invention comprises Fe, Zr, Nb and B in addition to Zn, and at least 50% or more of the structure It consists of microcrystal grains of bccFe with an average grain size of 100 nm or less, and the balance consists of an amorphous alloy.

In order to solve the above problems, the third highly saturated magnetic flux density low iron loss Fe-based soft magnetic alloy of the present invention is the iron loss by heating at 320 ° C. for 100 hours in the highly saturated magnetic flux density low iron loss Fe-based magnetic alloy described above. The change rate is 20% or less, the saturation magnetic flux density is 1.5 T or more, and the magnetic permeability is 30,000 or more.

In order to solve the above problems, the third highly saturated magnetic flux density low iron loss Fe-based soft magnetic synthesis of the present invention is composed of a composition represented by the following formula, and at least 50% or more of the structure has a fine grain size of bccFe of 100 nm or less It consists of crystal grains, the balance is made of amorphous alloy phase, the microcrystalline grain of bccFe is to quench the alloy to form a nearly amorphous single-phase structure, and then the amorphous phase is cooled and precipitated after heating above the crystallization temperature.

(Fe _1-a Q _a ) _b Zr _x Nb _y B _t Zn _z

However, Q is either Co or Ni, or both, and a, b, x, y, t, and z, which represent the composition ratio, are 0. a 0.05, 80 atomic% b, 1.5 atomic% x 2.5 atomic%, 3.5 atomic% y 5.0 atomic%, 5 atomic% t 12.5 atomic%, 0.025 atomic% z 0.2 atomic%, 5.0 atomic% x + y 7.5 atomic%, 1.5 / 6 x / (x + y) 2.5 / 6.

In order to solve the above problems, the third high saturation magnetic flux density low iron loss Fe-based magnetic alloy of the present invention is composed of a composition represented by the following formula, wherein at least 50% or more of the structure has a fine grain size of bccFe of 100 nm or less It consists of crystal grains, the balance consists of amorphous alloy phase, the microcrystalline grain of bccFe is to quench the alloy to form a nearly amorphous single-phase structure, and then the amorphous phase is cooled and precipitated after heating above the crystallization temperature.

(Fe _1-a Q _a ) _b Zr _x Nb _y B _t Zn _z M ' _u

However, Q is either or both of Co and Ni, and M 'is one or two or more elements selected from Cr, Ru, Rh and Ir, and represents a, b, x, y, t, z indicating composition ratio. , u is 0 a 0.05, 80 atomic% b, 1.5 atomic% x 2.5 atomic%, 3.5 atomic% y 5.0 atomic%, 5 atomic% t 12.5 atomic%, 0.025 atomic% z 0.2 atomic%, u 5 atomic%, 5.0 atomic% x + y 7.5 atomic%, 1.5 / 6 x / (x + y) 2.5 / 6.

The high saturation magnetic flux density low iron loss Fe-based soft magnetic alloy of the present invention has a fracture strain of 10 × 10 ⁻³ or more in order to solve the above problems.

The low iron loss magnetic core of the present invention uses the above-described Fe-based soft magnetic alloy, and the low iron loss magnetic core of the present invention comprises a low iron-loss Fe-based soft magnetic alloy having a saturation magnetic flux density of 1.5 T or more.

Further, the low iron loss magnetic core of the present invention consists of a low iron loss Fe-based magnetic alloy having an iron loss of 0.15 W / kg or less when a 1.4 T magnetic flux is applied at a frequency of 50 Hz.

Further, the low iron loss magnetic core of the present invention consists of a low iron loss Fe-based magnetic alloy having a rate of change of iron loss of 10% or less before and after aging at 200 ° C for 500 hours.

The low iron loss magnetic core of the present invention is composed of a ribbon of low iron loss Fe-based soft magnetic alloy having a fracture strain of 1.0 × 10 ⁻² or more.

The low iron loss magnetic core of the present invention is obtained by laminating one or two or more annular bodies formed of a ribbon of the low iron loss Fe-based soft magnetic alloy.

In addition, the low iron loss magnetic core of the present invention is composed of a ring in which the low iron loss Fe-based soft magnetic alloy is wound in a ring shape.

In order to achieve the above object, the manufacturing method of the Fe-based soft magnetic alloy in the present invention adopts the following configuration.

In the first production method of the present invention, the temperature increase rate until reaching the heat treatment temperature in the heat treatment step is set to 10 ° C / min or more and 200 ° C / min or less for various alloys as shown below.

Moreover, More preferably, the said temperature increase rate is set to 30 degreeC / min or more and 100 degrees C / min or less.

In addition to the above, the heat treatment temperature is set to 490 ° C or higher and 670 ° C or lower, and more preferably 500 ° C or higher and 560 ° C or lower.

The second manufacturing method of the Fe-based soft magnetic alloy according to the present invention is 1 selected from Ti, Zr, Hf, V, Nb, Ta, Mo, W, and Mn based on the above-mentioned Fe melted in a crucible. A method for producing a Fe-based soft magnetic alloy in which an alloy melt containing species or two or more elements M and B is injected into a cooling roll by a nozzle, and rapidly quenched and solidified on the cooling roll, wherein the melt is injected from the nozzle. The temperature is lower than 1350 ° C.

In this method, the alloy, which is in an amorphous state of rapid solidification, is heated above the crystallization temperature to precipitate bccFe crystal grains. Then, the injection temperature of the melt is 1240 ℃ or more.

The Fe-based soft magnetic alloy is composed of microcrystal grains mainly composed of bccFe having an average grain size of 100 nm or less and at least 50% or more of the structure, and the balance having an amorphous structure.

In order to solve the above problems, the method for producing the third Fe-based soft magnetic alloy of the present invention includes Fe as a main component, and is selected from Ti, Zr, Hf, V, Nb, Ta, Mo, W, Mn, or An amorphous alloy containing two or more kinds of elements M and B was subjected to a first heat treatment to form a microcrystalline alloy containing an amorphous phase mainly composed of fine grains of Fe having a fine bcc structure having an average grain size of 30 nm or less, followed by 100 캜 or more. The second heat treatment is performed at a holding temperature below the holding temperature of the first heat treatment temperature.

In addition, the method for producing a third Fe-based soft magnetic alloy of the present invention, in order to solve the above problems, in the manufacturing method of the Fe-based soft magnetic alloy described above, the holding temperature of the second heat treatment to 200 to 400 ℃ It is.

In order to solve the above problems, the method of manufacturing the third Fe-based soft magnetic alloy according to the present invention is to maintain the second heat treatment for 0.5 to 100 hours in the method of manufacturing the Fe-based soft magnetic alloy described above. .

In addition, the manufacturing method of the third Fe-based soft magnetic alloy of the present invention is performed by maintaining the second heat treatment for 1 to 30 hours in the manufacturing method of the Fe-based soft magnetic alloy described above in order to solve the above problems. will be.

In addition, the method for producing a third Fe-based soft magnetic alloy of the present invention, in order to solve the above problems, in the method for producing a Fe-based soft magnetic alloy described above, the first heat treatment is a temperature increase of 10 to 20 ℃ / min It is done at speed.

And in order to solve the above problems, the manufacturing method of the Fe-based soft magnetic alloy of the present invention, in the manufacturing method of the Fe-based soft magnetic alloy described above, characterized in that the holding temperature of the first heat treatment is 500 to 800 ℃ do.

And the manufacturing method of the 1st-3rd Fe-based soft magnetic alloy of this invention mentioned above, It is more preferable to use what is represented by the following composition formula as said Fe-based soft magnetic alloy.

(Fe _1-a Z _a ) _b B _x M _y

Provided that Z is one or two or more elements of Ni and Co, M is one or two or more elements selected from Ti, Zr, Hf, V, Nb, Ta, Mo, W, and Mn, and a 0.1, 75 atomic% b 93 atomic%, 0.5 atomic% x 18 atomic%, 4 atomic% y 9 atomic percent.

In addition, the Fe-based soft magnetic alloy may be represented by the following formula.

(Fe _1-a Z _a ) _b B _x M _y X _Z

Provided that Z is one or two or more elements of Ni and Co, M is one or two or more elements selected from Ti, Zr, Hf, V, Nb, Ta, Mo, W, and Mn, and X is Si, Al , Ge, Ga, and 0 a 0.1, 75 atomic% b 93 atomic%, 0.5 atomic% x 18 atomic%, 4 atomic% y 9 atomic%, z 5 atomic percent.

In addition, the Fe-based soft magnetic alloy may be represented by the following formula.

(Fe _1-a Z _a ) _b B _x M _y T _t

Provided that Z is one or two or more elements of Ni and Co, M is one or two or more elements selected from Ti, Zr, Hf, V, Nb, Ta, Mo, W, and Mn, and T is Cu, Ag , Au, Pd, Pt is one or two or more elements selected from 0, a 0.1, 75 atomic% b 93 atomic%, 0.5 atomic% x 18 atomic%, 4 atomic% y 9 atomic%, t 5 atomic percent.

In addition, the Fe-based soft magnetic alloy may be represented by the following formula.

(Fe _1-a Z _a ) _b B _x M _y T _t X _z

Provided that Z is one or two or more elements of Ni and Co, M is one or two or more elements selected from Ti, Zr, Hf, V, Nb, Ta, Mo, W, and Mn, and T is Cu, Ag , One or two or more elements selected from Au, Pd, and Pt, X is Si, Al, Ge, Ga, and 0 a 0.1, 75 atomic% b 93 atomic%, 0.5 atomic% x 18 atomic%, 4 atomic% y 9 atomic%, t 5 atomic%, z 5 atomic percent.

In addition, the Fe-based soft magnetic alloy may use the first to third highly saturated magnetic flux density low iron loss Fe-based soft magnetic alloy.

BRIEF DESCRIPTION OF THE DRAWINGS It is the block diagram which made a cross section part of an example of the apparatus for manufacturing the alloy of this invention.

It is a figure which shows the relationship between the amount of Zn put into the crucible used when manufacturing the alloy of this invention, and the Zn analysis value of the obtained alloy ribbon sample.

3 is an exploded perspective view showing a low iron loss magnetic core according to an embodiment of the present invention.

4 is an exploded perspective view showing another low iron loss magnetic core of an embodiment of the present invention.

5 is a perspective view showing a common mode choke coil using the low iron loss magnetic core of the present invention.

6 is a graph showing an example of a heat treatment pattern of the method of manufacturing a third Fe-based soft magnetic alloy of the present invention.

7 is a graph showing an example of a heat treatment pattern of the method of manufacturing a third Fe-based soft magnetic alloy of the present invention.

Fig. 8 shows the results of X-ray diffraction analysis before heat treatment of a highly saturated magnetic flux density low iron loss Fe-based soft magnetic alloy having a composition of Fe _85.75 Zr ₂ Nb ₄ B _8.25 .

Fig. 9 shows the results of X-ray diffraction analysis before heat treatment of a highly saturated magnetic flux density low iron loss Fe-based soft magnetic alloy having a composition of Fe _85.75 Zr ₂ Nb ₄ B _8.25 .

Fig. 10 is a graph showing the relationship between the composition of the soft magnetic alloy containing Zr and Nb in an amount of 5 atomic percent and the coercive force (Hc).

Fig. 11 is a graph showing the relationship between the composition of a soft magnetic alloy containing a total amount of Zr and Nb of 5 atomic% and the magnetic permeability (μ ') at 1 ㎑.

Fig. 12 is a graph showing the relationship between the composition of a soft magnetic alloy in which the total amount of Zr and Nb is 5 atomic% and the saturation magnetic flux density (B ₁₀ ) in the applied magnetic field of 10 Oe.

Fig. 13 is a graph showing the relationship between the composition of the soft magnetic alloy containing Zr and Nb in an amount of 5 atomic percent and the residual magnetization (Br).

Fig. 14 is a graph showing the relationship between the composition of the soft magnetic alloy containing Zr and Nb in an amount of 5.5 atomic% and the coercive force (Hc).

Fig. 15 is a graph showing the relationship between the composition of a soft magnetic alloy in which the total amount of Zr and Nb is 5.5 atomic% and the permeability (μ ').

Fig. 16 is a graph showing the relationship between the composition of a soft magnetic alloy in which the total amount of Zr and Nb is 5.5 atomic% and the saturation magnetic flux density (B ₁₀ ) in the applied magnetic field of 10 Oe.

Fig. 17 is a graph showing the relationship between the composition of the soft magnetic alloy containing Zr and Nb in an amount of 5.5 atomic percent and the residual magnetization (Br).

Fig. 18 is a graph showing the relationship between the composition of a soft magnetic alloy in which the total amount of Zr and Nb is 5.5 atomic% and the heat treatment temperature Ta when the coercive force Hc is minimum.

Fig. 19 is a graph showing the relationship between the composition of a soft magnetic alloy in which the total amount of Zr and Nb is 5.5 atomic% and the heat treatment temperature Ta when the magnetic permeability μ 'is maximized.

Fig. 20 is a graph showing the relationship between the composition of a soft magnetic alloy in which the total amount of Zr and Nb is 5.5 atomic% and the crystallization temperature (T _x1 ) of the bcc-Fe phase.

Fig. 21 is a graph showing the relationship between the composition of a soft magnetic alloy in which the total amount of Zr and Nb is 5.5 atomic% and the crystallization temperature (T _x1 ) of the compound phase precipitated on the high temperature side on bcc-Fe.

Fig. 22 is a graph showing the relationship between the composition of a soft magnetic alloy in which the total amount of Zr and Nb is 5.5 atomic% and the crystallization temperature (T _x2 ) of the FeB _x phase.

Fig. 23 is a graph showing the relationship between the composition of a soft magnetic alloy in which the total amount of Zr and Nb is 5.5 atomic% and the interval (ΔT _x ) between the crystallization temperatures;

Fig. 24 is a graph showing the relationship between the composition of a soft magnetic alloy containing zirconium and niobium in an amount of 6 atomic% and the coercive force (Hc).

Fig. 25 is a graph showing the relationship between the composition of a soft magnetic alloy containing zirconium and niobium in a total amount of 6 atomic percent and the permeability (μ ');

Fig. 26 is a graph showing the relationship between the composition of a soft magnetic alloy in which the total amount of Zr and Nb is 6 atomic% and the saturation magnetic flux density (B ₁₀ ) in the applied magnetic field of 10 Oe;

Fig. 27 is a graph showing the relationship between the composition of a soft magnetic alloy containing zirconium and niobium in an amount of 6 atomic% and the residual magnetization (Br);

Fig. 28 shows the relationship between the composition of a soft magnetic alloy containing zirconium and niobium in an amount of 6 atomic% and the average grain size of the bcc-Fe phase;

Fig. 29 is a graph showing the relationship between the composition of a soft magnetic alloy containing zirconium and niobium in an amount of 6 atomic% and the magnetostriction constant (λ _s );

Fig. 30 is a graph showing the relationship between the composition of a soft magnetic alloy containing a total amount of Zr and Nb of 6 atomic% and the heat treatment temperature Ta when the coercive force Hc is minimum.

FIG. 31 is a graph showing the relationship between the composition of a soft magnetic alloy containing zirconium and niobium in an amount of 6 atomic% and the heat treatment temperature Ta when the magnetic permeability μ 'is maximized;

Fig. 32 is a graph showing the relationship between the composition of a soft magnetic alloy in which the total amount of Zr and Nb is 6 atomic% and the crystallization temperature (T _x1 ) of the bcc-Fe phase.

Fig. 33 shows the relationship between the composition of a soft magnetic alloy containing zirconium and niobium in an amount of 6 atomic% and the crystallization temperature (T _x2 ) of the FeB _x phase.

Fig. 34 is a graph showing the relationship between the composition of a soft magnetic alloy containing zirconium and niobium in an amount of 6 atomic% and the crystallization temperature (T _x1 ) of a compound phase deposited on the high temperature side of bcc-Fe;

Fig. 35 is a graph showing the relationship between the composition of a soft magnetic alloy containing zirconium and niobium in a total amount of 6 atomic percent and the interval (? T _x ) between crystallization temperatures;

Fig. 36 is a graph showing the relationship between the composition of a soft magnetic alloy containing Zr and Nb in an amount of 6.5 atomic% and the coercive force (Hc).

Fig. 37 shows the relationship between the composition of a soft magnetic alloy containing zirconium and niobium in an amount of 6.5 atomic percent and the magnetic permeability (μ ').

Fig. 38 is a graph showing the relationship between the composition of a soft magnetic alloy in which the total amount of Zr and Nb is 6.5 atomic% and the saturation magnetic flux density (B ₁₀ ) in the applied magnetic field of 10 Oe;

Fig. 39 is a graph showing the relationship between the composition of a soft magnetic alloy containing zirconium and niobium in an amount of 6.5 atomic% and the residual magnetization (Br);

Fig. 40 is a graph showing the relationship between the composition of a soft magnetic alloy containing zirconium and niobium in an amount of 6.5 atomic% and the average grain size of the bcc-Fe phase;

Fig. 41 is a graph showing the relationship between the composition of a soft magnetic alloy containing zirconium and niobium in an amount of 6.5 atomic% and the magnetostriction constant (λ _s );

Fig. 42 is a graph showing the relationship between the composition of a soft magnetic alloy in which the total amount of Zr and Nb is 6.5 atomic% and the heat treatment temperature Ta when the coercive force Hc is minimum.

Fig. 43 is a graph showing the relationship between the composition of a soft magnetic alloy containing zirconium and niobium in an amount of 6.5 atomic% and the heat treatment temperature Ta when the magnetic permeability µ 'is maximum;

Fig. 44 shows the relationship between the composition of a soft magnetic alloy containing zirconium and niobium in an amount of 6.5 atomic% and the crystallization temperature (T _x1 ) of the bcc-Fe phase.

Fig. 45 is a graph showing the relationship between the composition of a soft magnetic alloy containing zirconium and niobium in an amount of 6.5 atomic% and the crystallization temperature (T _{x1 '} ) of a compound phase precipitated on the high temperature side of bcc-Fe.

Fig. 46 is a graph showing the relationship between the composition of a soft magnetic alloy in which the total amount of Zr and Nb is 6.5 atomic% and the crystallization temperature (T _x2 ) of the FeB _x phase.

Fig. 47 is a graph showing the relationship between the composition of a soft magnetic alloy in which the total amount of Zr and Nb is 6.5 atomic% and the interval (ΔT _x ) of the crystallization temperature.

Fig. 48 is a graph showing the relationship between the composition of a soft magnetic alloy containing zirconium and niobium in a total amount of 7 atomic percent and the coercive force (Hc).

Fig. 49 is a graph showing the relationship between the composition of a soft magnetic alloy containing zirconium and niobium in a total amount of 7 atomic percent and the magnetic permeability (μ ');

Fig. 50 is a graph showing the relationship between the composition of a soft magnetic alloy in which the total amount of Zr and Nb is 7 atomic percent and the saturation magnetic flux density (B ₁₀ ) in the applied magnetic field of 10 Oe;

Fig. 51 is a graph showing the relationship between the composition of a soft magnetic alloy containing zirconium and niobium in a total amount of 7 atomic percent and the residual magnetization (Br);

Fig. 52 is a graph showing the relationship between the composition of a soft magnetic alloy containing zirconium and niobium in a total amount of 7 atomic percent and the average grain size of the bcc-Fe phase;

Fig. 53 is a graph showing the relationship between the composition of a soft magnetic alloy containing zirconium and niobium in a total amount of 7 atomic percent and the magnetostriction constant (λ _s ).

Fig. 54 shows the relationship between the composition of a soft magnetic alloy containing zirconium and niobium in an amount of 7 atomic percent and the heat treatment temperature Ta when the coercive force Hc is minimum;

Fig. 55 is a graph showing the relationship between the composition of a soft magnetic alloy containing zirconium and niobium in a total amount of 7 atomic percent and the heat treatment temperature Ta when the magnetic permeability µ 'is maximum;

Fig. 56 is a graph showing the relationship between the composition of a soft magnetic alloy containing zirconium and niobium in an amount of 7 atomic% and the crystallization temperature (T _x1 ) of the bcc-Fe phase;

Fig. 57 is a graph showing the relationship between the composition of a soft magnetic alloy containing zirconium and niobium in an amount of 7 atomic% and the crystallization temperature (T _x2 ) of the FeB _x phase;

Fig. 58 is a graph showing the relationship between the composition of a soft magnetic alloy containing zirconium and niobium in a total amount of 7 atomic percent and the crystallization temperature (T _x1 ) of a compound phase precipitated on the high temperature side of bcc-Fe;

Fig. 59 is a graph showing the relationship between the composition of a soft magnetic alloy containing zirconium and niobium in an amount of 7 atomic% and the spacing (? T _x ) between crystallization temperatures;

Fig. 60 is a graph showing the relationship between the composition ratio of Zr and Nb and the coercive force (Hc).

Fig. 61 shows the relationship between the composition ratio of Zr and Nb and the interval (ΔT _x ) between the crystallization temperatures.

FIG. 62 shows the relationship between the heat treatment temperature (Ta) and fracture strain (λ _f ) of a soft magnetic alloy composed of Fe _85.5 Zr ₂ Nb ₄ B _8.5 , Fe ₉₀ Zr ₇ B ₃ , and Fe ₈₄ Nb ₇ B ₉ Drawing.

63 shows magnetic flux density (Bm) of a soft magnetic alloy composed of Fe ₇₈ Si ₉ B ₁₃ , Fe ₈₅ Zr _1.75 Nb _4.25 B ₉ , Fe _85.5 Zr ₂ Nb ₄ B _8.5 , Fe _85.75 Zr _2.25 Nb _3.75 B _8.25 It is a figure which shows the relationship of iron loss (Bm).

Fig. 64 is a graph showing the relationship between magnetic flux density (Bm) and iron loss of a soft magnetic alloy composed of Fe ₇₈ Si ₉ B ₁₃ and Fe ₈₅ Zr _1.75 Nb _4.25 B ₉ .

FIG. 65 is a graph showing the time-lapse change of iron loss of a soft magnetic alloy composed of Fe ₇₈ Si ₉ B ₁₃ , Fe ₈₅ Zr _1.75 Nb _4.25 B ₉ , Fe _85.75 Zr _2.25 Nb _3.75 B _8.25 .

66 shows magnetic flux density (Bm) of a soft magnetic alloy composed of Fe ₇₈ Si ₉ B ₁₃ , Fe ₈₅ Zr _1.75 Nb _4.25 B ₉ , Fe _85.5 Zr ₂ Nb ₄ B _8.5 , Fe _85.75 Zr _2.25 Nb _3.75 B _8.25 It is a figure which shows the relationship of iron loss change rate.

FIG. 67 is a graph showing the time-lapse change of iron loss of a soft magnetic alloy composed of Fe ₇₈ Si ₉ B ₁₃ and Fe _85.5 Zr ₂ Nb ₄ B _8.5 .

FIG. 68 is a view showing a time-lapse change in iron loss rate of a soft magnetic alloy composed of Fe ₇₈ Si ₉ B ₁₃ and Fe _85.5 Zr ₂ Nb ₄ B _8.5 .

Fig. 69 is an X-ray diffraction _pattern before and after heat treatment of a sample having a composition of (Fe _0.8575 Zr _0.02 Nb _0.04 B _0.0825 ) _99.88 Zn _0.12 of the composition of the present invention.

70 shows the results of measuring the coercive force of a sample having a composition composed of Fe _c Zr _d Nb _e B _f of a composition system similar to the composition of the present invention and (Fe) of the composition system in which Zn was added in the range of 0.034 to 0.142 atomic%. _c Zr _d Nb _e B _f ) A triangular composition diagram showing the results of measuring the coercive force of a sample of composition consisting of _100-z Zn _z .

FIG. 71 is a triangular composition diagram showing the result of measuring the magnetic permeability (μ ': real part of the magnetic permeability) at 1 kHz with respect to a sample having the same composition as the sample shown in FIG.

FIG. 72 is a triangular composition diagram showing the saturation magnetic flux density (B ₁₀ ) obtained from a magnetization curve obtained by applying an applied magnetic field of 10 Oe to a sample having the same composition as the sample shown in FIG. 71.

Fig. 73 is a triangular composition diagram showing a measurement result of residual magnetic flux density (Br) of the previous sample.

Fig. 74 is a triangular composition diagram showing the first crystallization temperature (T _x1 is the crystallization temperature of bccFe) of the previous sample.

Fig. 75 is a triangular composition diagram showing the intermediate crystallization temperature (T _x1 'is the crystallization temperature of the compound phase) of the previous sample.

Fig. 76 is a triangular composition diagram showing the second crystallization temperature (T _x2 is the crystallization temperature of the compound phase) of the previous sample.

Fig. 77 is a triangular composition diagram showing ΔT _x expressed by T _x2 -T _x1 in the previous sample.

Fig. 78 is a triangular composition diagram showing the grain size of a sample of the composition system not containing Zn as a composition system similar to the composition of the present invention.

Fig. 79 is a triangular composition diagram showing the magnetostriction (? S) of a sample of a composition system not containing Zn as a composition system similar to the composition of the present invention.

Fig. 80 shows the Zn concentration dependence of the grain size (D) of the composition-based alloy sample of the present invention with Zn added.

Fig. 81 shows the Zn concentration dependence of the magnetostriction (? S) of the composition-based alloy sample of the present invention with Zn added.

82 is example compares the results of measuring the iron loss for the samples by the addition of Zn 0.12 at% or 0.13 atomic percent of the alloy samples of the composition consisting of Fe _85.75 Zr ₂ Nb ₄ B _8.25 to AC magnetization characteristic measuring apparatus Fe ₇₈ Si ₉ B ₁₃ is a view showing compared to the value of the ribbon sample of the composition.

83 is a comparative example sample of the composition consisting of Fe ₇₈ Si ₉ B ₁₃ , a comparative example sample of the composition consisting of Fe _85.75 Zr _2.25 Nb _3.75 B _8.25 , and (Fe _0.8575 Zr _0.02 Nb _0.04 B _0.0825 ) _99.88 Zn _0.12 It is a figure which shows the time-lapse change (measured at normal temperature after 200 degreeC heating) of the iron loss of the alloy sample of this invention of composition.

FIG. 84 is a view showing iron loss at room temperature after heating for 320 ° C. for a predetermined time using a sample having the same composition as the sample used in FIG. 17.

FIG. 85 is a diagram showing a rate of change of iron loss shown in FIG. 18.

Fig. 86 is a ribbon sample having a sheet thickness of 20 µm, respectively, showing values of bending diameter and fracture strain for the samples of the comparative examples of the various compositions and the alloy samples of the present invention.

FIG. 87 is a graph showing Zn concentration dependence on the Curie temperature change of the amorphous phase determined by the temperature change of magnetization. FIG.

Fig. 88 is a graph showing the Zn concentration dependence on the Curie temperature change of the ribbon sample obtained by the temperature change of magnetization.

Fig. 89 shows the Zn concentration dependence of the coercive force of the FeNbB alloy.

Fig. 90 is a graph showing the Zn concentration dependence of the permeability of FeNbB alloy.

91 is a graph showing the Zn concentration dependency of the coercive force of the FeZrNbB-based alloy.

Fig. 92 is a graph showing the Zn concentration dependence of the permeability of FeZrNbB alloy.

Fig. 93 is a triangular compositional diagram showing the coercive force of a sample containing 4 atomic% in total of Zr and Nb.

Fig. 94 shows the heat treatment temperature Ta and the temperature increase rate and the permeability of the alloy for the heat treatment in which the holding time after reaching the heat treatment temperature Ta was 0 min for the alloy composed of the composition Fe _85.5 Zr ₂ Nb ₄ B _8.5 . It is a graph showing the relationship with (μ ').

FIG. 95 shows the heat treatment temperature Ta and the temperature increase rate and the coercive force of the alloy with respect to an alloy having a composition Fe _85.5 Zr ₂ Nb ₄ B _8.5 having been subjected to a heat treatment in which the holding time after reaching the heat treatment temperature Ta was 0 min. It is a graph which shows the relationship of (Hc).

FIG. 96 is a graph showing the magnetic permeability μ ′ when the holding time is 5 min in FIG. 94.

FIG. 97 is a graph of the coercive force (Hc) when the holding time is 5 min in FIG. 95.

FIG. 98 is a graph of the magnetic permeability μ ′ when the holding time is 10 min in FIG. 94.

FIG. 99 is a graph of the coercive force (Hc) when the holding time is 10 min in FIG. 95.

FIG. 100 is a graph of the magnetic permeability μ ′ when the holding time is 30 min in FIG. 94.

FIG. 101 is a graph of the coercive force (Hc) when the holding time is 30 min in FIG. 95.

FIG. 102 is a graph of the magnetic permeability μ 'when the holding time is 60 min in FIG.

FIG. 103 is a graph of the coercive force (Hc) when the holding time is 60 min in FIG. 95.

Fig. 104 is a graph showing the magnetic permeability μ 'shown in Figs. 94, 96, 98, 100, and 102 at an injection temperature of 1280 ° C in one drawing.

Fig. 105 is a graph showing the coercive force Hc shown in Figs. 94, 96, 98, 100 and 102 at an injection temperature of 1280 ° C in one drawing.

FIG. 106 is a graph showing only the magnetic permeability (μ ′) for the holding time of 0, 10 and 60 min in FIG. 104.

FIG. 107 is a graph showing only the coercive force (Hc) for the holding time of 0, 10 and 60 min in FIG. 105.

FIG. 108 is a graph showing only the permeability (μ ′) for the holding time of 5 and 30 min in FIG. 104.

FIG. 109 is a graph showing only the coercive force Hc of the holding time of 5 and 30 min in FIG. 105.

FIG. 110 is a graph showing the magnetic permeability μ 'shown in FIGS. 95, 97, 99, 101 and 103 at the injection temperature of 1320 ° C in one view.

Fig. 111 is a graph showing the coercive force Hc shown in Figs. 95, 97, 99, 101 and 103 at the injection temperature of 1320 캜 in one drawing.

FIG. 112 is a graph showing only the magnetic permeability (μ ') for the holding time of 0, 10, and 60 min in FIG.

FIG. 113 is a graph showing only the coercive force (Hc) for the holding time of 0, 10 and 60 min in FIG. 111.

FIG. 114 is a graph showing only the permeability (μ ′) for the holding time of 5 and 30 min in FIG. 110.

FIG. 115 is a graph showing only the coercive force Hc of the holding time of 5 and 30 min in FIG. 111.

FIG. 116 shows the heat treatment temperature (Ta) and the temperature increase rate of the alloy having the composition Fe _85.5 Zr ₂ Nb ₄ B _8.5 having a holding time of 5 min after reaching the heat treatment temperature Ta; It is a graph showing the relationship with the magnetic permeability (μ ').

FIG. 117 shows the heat treatment temperature (Ta) and the temperature increase rate of the alloy having the composition Fe _85.5 Zr ₂ Nb ₄ B _8.5 having a holding time after reaching the heat treatment temperature Ta of 5 min. It is a graph showing the relationship between the coercive force (Hc).

118 is a graph showing the magnetic permeability (μ ′) when the composition (Fe _85.5 Zr ₂ Nb ₄ B _8.5 ) ₉₉ Zn ₁ is set in FIG. 116.

FIG. 119 is a graph showing the coercive force Hc when the composition (Fe _85.5 Zr ₂ Nb ₄ B _8.5 ) ₉₉ Zn ₁ is set in FIG. 117.

FIG. 120 is a graph showing the magnetic permeability (μ ′) when the composition (Fe _85.5 Zr ₂ Nb ₄ B _8.5 ) ₉₉ Zn ₂ is set in FIG. 116.

FIG. 121 is a graph showing the coercive force (Hc) when the composition (Fe _85.5 Zr ₂ Nb ₄ B _8.5 ) ₉₉ Zn ₂ is set in FIG. 117.

FIG. 122 is a graph showing the magnetic permeability (μ ′) when the composition (Fe _85.5 Zr ₂ Nb ₄ B _8.25 ) ₉₉ Zn ₃ is set in FIG. 116.

FIG. 123 is a graph showing the coercive force (Hc) when the composition (Fe _85.5 Zr ₂ Nb ₄ B _8.5 ) ₉₉ Zn ₃ is set in FIG. 117.

Fig. 124 is a graph showing the magnetic permeability μ 'shown in Figs. 116, 118, 120 and 122 at an injection temperature of 1260 deg.

FIG. 125 is a graph showing the coercive force Hc shown in FIGS. 117, 119, 121 and 123 at an injection temperature of 1260 ° C in one view.

FIG. 126 is a graph showing the magnetic permeability μ 'shown in FIGS. 116, 118, 120 and 122 at an injection temperature of 1300 deg.

FIG. 127 is a graph showing the coercive force Hc shown in FIGS. 117, 119, 121 and 123 at an injection temperature of 1300 ° C in one view.

FIG. 128 is a view showing a second heat treatment time dependency of the coercive force Hc of the soft magnetic alloys of various compositions.

129 is a view showing the second heat treatment time dependency of the magnetic permeability of soft magnetic alloys of various compositions.

EMBODIMENT OF THE INVENTION Hereinafter, embodiment of this invention is described in detail.

First, the first highly saturated magnetic flux density low iron loss Fe-based soft magnetic alloy of the present invention will be described.

In the first highly saturated magnetic flux density low iron loss Fe-based soft magnetic alloy of the present invention, the alloy obtained by the process obtained by quenching an alloy mainly composed of the amorphous phase in the melt is melted to precipitate a fine crystalline structure composed of fine crystal grains. It can obtain normally by the heat processing to make it.

The high saturation magnetic flux density low iron loss Fe-based soft magnetic alloy of the composition is a bcc-Fe phase (body core) having an average grain size of 100 nm or less, more preferably 30 nm or less, which accounts for 50% or more, preferably 70% or more of the structure. It is composed mainly of fine crystalline structure consisting of fine crystal grains mainly composed of a cubic Fe phase) and the remaining amorphous structure, and thus exhibits low magnetostriction, high saturation magnetic flux density, and excellent permeability.

In the composition system of the present invention, Fe, which is a main component, is an element that is responsible for magnetism and is important for obtaining high saturation magnetic flux density and excellent soft magnetic properties.

In the highly saturated magnetic flux density low iron loss Fe-based soft magnetic alloy of these compositions, a representing the composition ratio of Fe is 80 atom% or more. If a is less than 80 atomic%, the saturation magnetic flux density cannot be 1.5 T or more, and therefore is not preferable.

Moreover, it is more preferable that the composition ratio a of Fe is 83 atomic% or more, and it is more preferable to set it as 85 atomic% or more and 86 atomic% or less.

In order to obtain a saturation magnetic flux density of 1.5 T or more, it is necessary to contain as much Fe as possible after satisfying the range of addition of other additive elements, and in consideration of the amount of other additive elements, it is described later by containing an amount exceeding 80 atomic%. As shown in the example, the saturation magnetic flux density of 1.5 T or more can be easily obtained. Moreover, when a is 85 atomic% or more, saturation magnetic flux density can be 1.6T or more.

Moreover, when a exceeds 86 atomic%, it becomes difficult to obtain an amorphous single phase by the liquid quenching method, and as a result, the structure of the alloy obtained after heat treatment becomes uneven and high permeability cannot be obtained. Therefore, it is not preferable.

Moreover, a part of Fe may be substituted by Co or Ni for adjustment of magnetostriction, etc. In this case, Preferably it is 10% of Fe, More preferably, it is 5% or less. Outside this range, the permeability deteriorates.

In the present invention, in order to easily obtain an amorphous phase, it is necessary to simultaneously contain Zr and Nb having high amorphous forming ability.

In addition, B is considered to have the effect of increasing the amorphous formability of the highly saturated magnetic flux density low iron loss Fe-based soft magnetic alloy of the present invention and controlling the formation of a compound phase which adversely affects the magnetic properties in the heat treatment. Addition is essential.

Zr and Nb are important elements for obtaining an amorphous phase when quenched in an alloy melt, and are important for achieving high saturation magnetic flux density and high permeability of 1.5 T or more by precipitating fine grains of bcc-Fe phase by heat treatment in the amorphous phase. Do.

When the composition ratio of Zr is x and the composition ratio of Nb is y, the required amount of amorphous phase cannot be obtained unless these elements are added at a total amount (x + y) of 5 atomic% or more. When the value of (x + y) exceeds 7 atomic%, the saturation magnetic flux density deteriorates and the soft magnetic properties deteriorate, which is not preferable.

Further, by setting the composition ratio of Zr and Nb in the range of 1.5 / 6 ≤ x / (x + y) ≤ 2.5 / 6, the magnetic permeability and saturation magnetic flux density of the high saturation magnetic flux density low iron loss Fe iron magnetic alloy are further improved and the iron loss This becomes very small. Especially in the low frequency region of several hundreds of kHz or less, the iron loss is remarkable, and the change over time of iron loss is also small.

The total amount of Zr and Nb is more preferably 5 atomic% or more and 7 atomic% or less, and still more preferably 5.7 atomic% or more and 6.5 atomic% or less. Moreover, it is more preferable to make the composition ratio of Zr and Nb into x / (x + y) = 2/6.

Moreover, in order to acquire favorable soft magnetic characteristics, the composition ratio x of Zr may be 0.5 atomic% or more and 3.5 atomic% or less, More preferably, you may be 1.5 atomic% or more and 2.5 atomic% or less. In order to obtain good soft magnetic properties, the composition ratio y of Nb may be 3 atomic% or more and 5.5 atomic% or less, more preferably 3.5 atomic% or more and 5.0 atomic% or less.

In addition, as described above, since B also has an amorphous forming ability, Zr and Nb contribute to amorphous formation, but when added more than necessary, a tendency of forming a compound phase which lowers the magnetic permeability and deteriorates soft magnetic properties between Fe and Since it becomes high, it is necessary to make composition ratio z into 5 atomic% or more and 12.5 atomic% or less, More preferably, it is 6 atomic% or more and 9.5 atomic% or less, More preferably, it is 8 atomic% or more and 9 atomic% or less.

In addition, although Zr and Nb are hardly dissolved in bcc-Fe, Zr and Nb are supersaturated in bcc-Fe by quenching the alloy to amorphous, and then performing heat treatment to crystallize. This heat treatment improves the soft magnetic properties of the high saturation magnetic flux density, low iron loss Fe-based soft magnetic alloys obtained by controlling the high capacity of these elements and depositing them into fine crystalline structures, thereby reducing the magnetostriction of the alloy ribbon.

In addition, in order to precipitate fine crystalline tissues and suppress coarsening of crystal grains in the fine crystalline tissues, it is considered that it is necessary to leave an amorphous phase at the grain boundaries, which may be an obstacle to grain growth.

In addition, this grain boundary amorphous phase is thought to suppress the formation of Fe-Zr-based or Fe-Nb-based compounds that deteriorate soft magnetic properties by retrieving Zr and Nb discharged from the bcc-Fe phase due to an increase in the heat treatment temperature. Therefore, it becomes important to add B to Fe-Zr system, Fe-Nb system, or Fe-Zr-Nb system alloy.

In addition to these elements, it is also possible to add platinum group elements such as Ru, Rh, Pd, Os, Ir, and Pt in order to improve the corrosion resistance. When these elements are added in an amount greater than 5 atomic%, the deterioration of the saturation magnetic flux density becomes remarkable, so the amount of addition must be 5 atomic% or less.

In addition, if necessary, Y, La, Ce, Pr, Nd, Pm, Sm, Eu, Gd, Tb, Dy, Ho, Er, Tm, Yb, Lu, Zn, Cd, In, Sn, Pb, As The magnetostriction of the highly saturated magnetic flux density low iron loss Fe-based soft magnetic alloy can also be adjusted by adding elements such as Sb, Bi, Se, Te, Li, Be, Mg, Ca, Sr, and Ba.

In addition, in the case of the highly saturated magnetic flux density low iron loss Fe-based soft magnetic alloy of the composition system, the undesired impurities such as H, N, O, and S are not deteriorated, and preferably 0.1 atomic% or less. Of course you may.

The high saturation magnetic flux density low iron loss Fe-based soft magnetic alloy of the present invention has the crystallization temperature of the bcc-Fe phase as T _X1 , the crystallization temperature of the compound phase crystallized at a higher temperature side than T _X1 as T _X2 , and the crystallization temperature. the interval _{(ΔT X) ΔT X = T} X2 - T _X1 when in, preferably in the 200 ℃ ≤ΔT _X.

Here, the other compound phase is a phase to be crystallized next to the bcc-Fe phase on the higher temperature side than T _X1 as described above, but the composition of the phase is considered to be Fe ₃ B or Fe ₂ B or the like.

When the alloy mainly composed of the amorphous phase generated by quenching is heated, exothermic reaction occurs by crystallization of the bcc-Fe phase, and the exothermic reaction is carried out by crystallization of another compound phase (Fe ₃ B or Fe ₂ B, etc.) at a predetermined temperature interval. Reaction takes place. These exothermic reactions are exothermic peaks on the DTA curve, for example, by performing differential thermal analysis (DTA measurement) on the alloy after quenching to obtain respective crystallization temperatures (T _X1 , T _X2 ) and the interval between the crystallization temperatures. (ΔT _X ) can also be obtained.

If the interval (ΔT _X ) of the crystallization temperature thus obtained is 200 ° C. or more, the interval between the crystallization temperature of the bcc-Fe phase and the compound phase becomes wider. Therefore, only the bcc-Fe phase is precipitated by heat treatment of the alloy under optimum conditions. It is possible to suppress the precipitation of other compound phases and to improve the soft magnetic properties of the highly saturated magnetic flux density low iron loss Fe-based soft magnetic alloy.

Specifically, the heat treatment temperature T _a of the alloy is preferably in the range of T _X1 <T _a <T _X2 .

Further, in the DTA measurement there is a case that other different precipitation in between the T _X1 and T _X2. Although the composition of this other phase is not clear, the precipitation of this other phase tends to be governed by the composition of the high saturation magnetic flux density low iron loss Fe-based soft magnetic alloy, and is particularly different from B because it is easily precipitated by increasing the composition ratio of B. It is considered to be a compound phase with an additive element, but it is not certain.

When the crystallization temperature of this other phase is set to T _X1 , the interval ΔT _X of the crystallization temperature can be set to T _X = T _X1 -T _X1 . In this case, it is, of course, preferable that the interval (ΔT _X ) of the crystallization temperature is not lower than 200 ° C.

In addition, when the other phase deposition, it is preferable that a heat treatment temperature of the alloy in _a range of T _X1 T <T _a <T _X1.

More specifically, the temperature rise temperature is 10 degrees C / min or more, More preferably, it is 10 degrees C / min or more and 200 degrees C / min or less, More preferably, it is 30 degrees C / min or more and 100 degrees C / min or less, and the heat processing temperature T _a is 490 degreeC. The heat treatment conditions are preferably 670 ° C. or less, more preferably 500 ° C. or more and 560 ° C. or less, and there is no holding time or 1 hour or less.

Moreover, it is preferable to make the heat processing atmosphere into a vacuum or inert gas atmosphere at the point which can prevent oxidation of an alloy.

The first highly saturated magnetic flux density low iron loss Fe-based soft magnetic alloy having the above-described configuration has a saturation magnetic flux density of 1.5 T or more and an iron loss of 0.15 W / kg when a magnetic flux of 1.4 T is applied at a frequency of 50 Hz. It becomes below and the iron loss change rate before and after ageing at 200 degreeC for 500 hours will be 10% or less. In addition, the highly saturated magnetic flux density low iron loss Fe-based soft magnetic alloy has a fracture strain of 1.0 × 10 ⁻² or more.

In the above-mentioned first highly saturated magnetic flux density low iron loss Fe-based soft magnetic alloy, 50% or more of the structure becomes a microcrystalline structure mainly composed of a bcc-Fe phase having an average grain size of 100 nm or less, and the composition shown in the following formula Since the magnetic permeability and the saturation magnetic flux density are improved, the iron loss can be reduced.

Fe _a Zr _x Nb _y B _z

Where a, x, y, and z represent composition ratios of 80 atomic% ≤ a, 5 atomic% ≤ x + y ≤ 7 atomic%, 1.5 / 6 ≤ x / (x + y) ≤ 2.5 / 6, 5 atomic% ≤ z ≤ 12.5 atomic%.

In the above-mentioned first highly saturated magnetic flux density low iron loss Fe-based magnetic alloy, the crystallization temperature of the bcc-Fe phase is T _X1 , and the crystallization temperature of the compound phase crystallized at a higher temperature side than the T _X1 side is T _X2 , When the interval ΔT _X of the crystallization temperature is ΔT _X = T _X2 -T _X1 , the crystallization temperature between the bcc-Fe phase and the compound phase crystallized on the high temperature side of the bcc-Fe phase is 200 ° C. ΔΔT _X. The spacing of the is widened, and by heat-treating the alloy under optimum conditions, only the bcc-Fe phase can be precipitated to suppress precipitation of other compound phases, thereby improving the soft magnetic properties of the soft magnetic alloy.

Next, the second highly saturated magnetic flux density low iron loss Fe-based soft magnetic alloy of the present invention will be described. One specific form of the soft magnetic alloy is composed of the composition shown in the following formula, wherein at least 50% or more of the tissue is composed of fine grains mainly composed of bcc-Fe having an average grain size of 100 nm or less, and the balance is in the form of amorphous alloys. The microcrystal grains mainly composed of bcc-Fe are formed by cooling an alloy (melt) having a composition described later to form a nearly amorphous single phase structure, and then cooling and depositing the amorphous phase after heating the crystallization temperature or more.

(Fe _1-a Q _a ) _b B _x M _y Zn _z

However, Q is either or both of Co and Ni, and M is one or two or more elements selected from Ti, Zr, Hf, V, Nb, Ta, Mo, and W, and a, b, x, y, z are 0 ≦ a ≦ 0.05, 80 atomic% ≦ b, 5 atomic% ≦ x ≦ 12.5 atomic%, 5 atomic% ≦ y ≦ 7 atomic%, 0.025 atomic% ≦ z ≦ 0.2 atomic%.

Another specific embodiment of the second soft magnetic alloy of the present invention has a composition formula in which the element Q is omitted in the above-described composition, ie, Fe _b B _x M _{y Z n z} . The ratio of each element in this composition formula is the same as that of the previous form.

Another specific form of the second soft magnetic alloy of the present invention consists of a composition represented by the following formula, wherein at least 50% or more of the tissue consists of microcrystal grains mainly composed of bcc-Fe having an average grain size of 100 nm or less, The microcrystalline grains composed mainly of bcc-Fe consisting of an amorphous alloy phase are quenched to form an almost amorphous single phase structure by quenching an alloy (melt) having a composition to be described later, and then cooling the amorphous phase after heating above the crystallization temperature. It is precipitated.

(Fe _1-a Q _a ) _b B _x M _y Zn _z M ' _u

Provided that Q is one or both of Co and Ni, M is one or more elements selected from Ti, Zr, Hf, V, Nb, Ta, Mo, and W, and M 'is Cr, Ru, Rh A, b, x, y, z, u, which represents one or more elements selected from among Ir and 0, a, b, x, y, z, u, are 0 ≦ a ≦ 0.05,80 atomic% ≦ b, 5 atomic% ≦ x ≦ 12.5 atomic%, 5 atomic% ≤ y ≤ 7 atomic%, 0.025 atomic% ≤ z ≤ 0.2 atomic%, u ≤ 5 atomic%.

The soft magnetic alloy of the above-mentioned composition is a step of obtaining an amorphous alloy ribbon or amorphous alloy powder containing an amorphous single phase or some crystals by quenching the alloy described later in a melt and a heat treatment for heating fine particles obtained by these steps to precipitate fine crystal grains. It can obtain normally by a process. However, in the alloy composition, since Zn tends to be easily lost by evaporation compared with other elements, it is necessary to set the amount of Zn to be added when producing a melt more than the range of the above-described composition formula as described in detail later. have.

In addition, what is obtained by said quenching method may be ribbon shape or powder shape, and of course, you may heat-process after shape-processing or machining to the shape aimed at obtained.

B is necessarily added to the soft magnetic alloy of the said composition. B has the effect of increasing the amorphous forming ability of the soft magnetic alloy, and the thermal stability of the Fe-M (= Zr, Hf, Nb, etc.)-Based microcrystalline alloy to become a barrier to grain growth, and has a thermally stable amorphous phase. There is an effect of remaining at the grain boundary. As a result, in the above-described heat treatment step, a fine body-centered cubic structure (bcc structure) having a particle size of 100 nm or less (specifically, 30 nm or less) that does not adversely affect magnetic properties under a wide heat treatment condition of 400 to 750 ° C. We can get organization mainly as subject. The content of B is preferably 5 atomic% or more and 12.5 atomic% or less, more preferably 6 atomic% or more and 9.5 atomic% or less, and most preferably 8 atomic% or more and 9.0 atomic% or less.

In addition, if necessary, Y, La, Ce, Pr, Nd, Pm, Sm, Eu, Gd, Tb, Dy, Ho, Er, Tm, Yb, Lu, Cd, In, Sn, Pb, As, Sb The magnetostriction of the soft magnetic alloy can be adjusted by adding elements such as, Bi, Se, Te, Li, Be, Mg, Ca, Sr, and Ba.

Moreover, since it does not interfere even if it contains in the range which does not reduce the characteristic of the alloy of this invention about the element which has amorphous formation ability, such as Al, Si, C, P, content of these elements should be about 1 atomic% or less. desirable. Moreover, as unavoidable impurity elements, such as H, N, O, and S, to the extent that the target characteristic does not deteriorate, Preferably it does not interfere even if it contains 0.1 atomic% or less.

In the soft magnetic alloy of the above composition, in order to easily obtain an amorphous phase with respect to Ti, Zr, Hf, V, Nb, Ta, Mo, and W represented by M, it is necessary to preferably contain Zr, Hf, and Nb having high amorphous forming ability. There is. In addition, Zr, Hf, and Nb can replace some of them with Ti, V, Ta, Mo, and W in other Periodic Tables 4A to 6A group elements.

Although Zr, Nb, and the like are originally hardly dissolved in bcc-Fe, Zr quenches the alloy melt by quenching and amorphous and then crystallizes by heat treatment. By adjusting the high capacity, the magnetostriction can be reduced. That is, the high capacity of Zr, Nb can be adjusted to the heat treatment conditions, and thus the value can be reduced by controlling the magnetostriction.

Therefore, in order to obtain low magnetostriction, it is necessary to obtain a fine crystal structure under a wide heat treatment condition. As described above, it is possible to obtain a fine crystal structure under a wide heat treatment condition by adding B, which results in both small magnetostriction and small magnetotropic anisotropy. It has a good magnetic characteristic as a result.

Therefore, in order to reduce iron loss in particular among the above-mentioned additive elements, it is particularly effective to set the ratio of Zr and Nb in a prescribed range. In the case where Zr and Nb are added as a main agent, the total amount thereof is 5? A relationship of Zr content + Nb content)? 7.5 is preferable, and a range of 5.7? (Zr content + Nb content)? 6.5 is more preferable.

Moreover, it is preferable to make the value of (Zr content) / (Zr content + Nb content) into the range of 1.5 / 6-2.5 / 6. This relationship is represented by the formula: 1.5 / 6 ≦ (Zr content) / (Zr content + Nb content) ≦ 2.5 / 6.

Moreover, (Zr content) / (Zr content + Nb content) = 2/6 is the most preferable among the range of this relational formula.

In addition, although the corrosion resistance is improved by adding Cr, Ru, Rh, and Ir to the composition as necessary, in order to maintain a high saturation magnetic flux density, it is preferable that the addition amount of these elements is 5 atomic% or less, and the saturation magnetic flux density and In consideration of both the soft magnetic properties and the iron loss, the content of 1 atomic% or less is more preferable.

The microcrystalline structure can be obtained by partially crystallizing the amorphous alloy of the Fe-M (= Zr, Hf) system by a special method, and the inventors of the present invention described in pages 217 to `` CONFERENCE ON METALLIC SCIENCE AND TECHNOLGY BUDAPEST '' in 1980. It is published on page 221. The same effect on the composition disclosed at this time was found in subsequent studies, and as a result, the present invention has been achieved. However, the reason for obtaining the microcrystal structure is already in the quench state of the amorphous phase forming step for producing the alloy of this system. It is thought that this is because a variation in the composition is generated, and this variation becomes a site of nonuniform nucleation, and a large number of uniform and fine nuclei are generated.

Content of Fe in the soft magnetic alloy of the said composition, or each content of Fe, Co, and Ni is 80 atomic% or more, Preferably it is less than 90 atomic%. This is because a high permeability cannot be obtained when these contents exceed 90 atomic%, but in order to obtain a saturation magnetic flux density of 1.55 T or more, the range of 83 to 87 atomic% (hereinafter, unless otherwise specified, the numerical ranges indicated by the upper limit is the upper limit). Since 83-87 atomic% shall mean 83 atomic% or more and 87 atomic% or less, since it is including the minimum and minimum, it is more preferable, and the range of 85-86 atomic% is still more preferable. Moreover, if Fe is not contained at least 80 atomic% or more, preferable saturation magnetic flux density will not be obtained.

Next, the Zn content in the soft magnetic alloy of the composition is preferably in the range of 0.025 atomic% or more and 0.2 atomic% or less. When added within this range, the coercive force and iron loss can be lowered and the permeability can be increased without lowering the high saturation magnetic flux density of 1.5 T or more.

In addition, the range of 0.034 atomic% or more and 0.16 atomic% or less is more preferable with respect to Zn content, and if it is this range, a soft magnetic alloy with low iron loss, a high saturation magnetic flux density, and little change with time can be obtained.

However, since Zn has a melting point of 419.5 ° C. and a boiling point of 908 ° C., when the soft magnetic alloy of the composition described above is melted in a crucible, the melting temperature is set to about 1240 to 1350 ° C., so most of Zn is evaporated and disappeared. do.

In order to rapidly quench the alloy melt of the above-mentioned composition into an amorphous alloy, the melt is injected into a cooling body such as a cooling roll to quench it, or an atomize method is performed to eject into the cooling gas. In order to be contained in the quenching alloy by the amount of, it is necessary to add Zn in an amount exceeding the amount of Zn in the alloy composition to be charged into the crucible.

That is, according to the researches of the present inventors, when using a melt at a temperature of about 1240 to 1350 ℃ to obtain an alloy such as a ribbon or powder in the melt by the quenching method, it is preferable to use a dosage of about 20 times or more of the target composition It turned out to be.

Fig. 1 shows an example of a manufacturing apparatus for obtaining a soft magnetic alloy of the present invention in a ribbon shape, and is connected to a vacuum pump 1 through an exhaust pipe 1a to enable vacuum exhaust inside a chamber 2. The copper or steel cooling roll 3 is provided so that rotation is free. The crucible 6 which has the nozzle 5 made of quartz etc. is provided in the upper part of this cooling roll 3, and the crucible 6 has a gas supply source 7 so that Ar gas pressure can be added to the crucible 6; ) Is connected via a gas supply pipe 7a, and a chamber 2 has a gas supply source 8 for adjusting the inside of the chamber 2 to a non-oxidizing gas pressure atmosphere such as Ar gas through the gas supply pipe 8a. Connected. In addition, an opening member is attached to the upper end of the crucible 6 so that the tip of the gas supply pipe 7a is connected to penetrate the opening member, and the inside of the crucible 6 is connected to the internal pressure of the chamber 2. It is configured to be pressurized separately.

A heating heater 9 is provided on the outer periphery of the bottom of the crucible 6, and is configured to heat-melt alloy raw materials introduced into the crucible 6 to obtain a melt, and the gas supply source 7 described above. By adding Ar gas pressure to the inside of the crucible 6, the melt is discharged to the surface of the cooling roll 3 during rotation through the nozzle 5, and is indicated by the reference numeral 11 of FIG. 1 on the side of the cooling roll 3. As described above, the ribbon is configured to be obtained.

Using the manufacturing apparatus shown in FIG. 1, the inside of the chamber 2 was made into an Ar gas atmosphere of about 160 Torr (Fe _0.94-t Zr _0.02 Nb _0.04 B _t ) _100-z Zn _z , t = 0.08, 0.0825, 0.085 Each ribbon sample in the case of preparing a ribbon by preparing each melt in which the Zn amount (z) was set to 1, 2, and 3 atomic% in the alloy having the composition to be sprayed, and ejecting the melt onto a rotating cooling roll. The result of having analyzed the Zn content of is shown in FIG.

As is clear from the results shown in Fig. 2, 0.035 to 0.0575 atomic% Zn remains in the ribbon sample obtained from the melt obtained by adding 1 atomic% Zn raw material into the crucible, and 2 atomic% Zn is added to the crucible (1). Zn of 0.07 to 0.125 atomic% remained in the ribbon sample obtained from the melt obtained, and Zn of 0.12 to 0.170 atomic% remained in the ribbon sample obtained from the melt obtained by injecting 3 atomic% Zn into the crucible (1). . In view of the above, it is understood that Zn in a ratio of 0.5 to 4.0 atomic% is required as the Zn raw material input amount at the time of forming the melt, in order to contain 0.025 to 0.2 atomic% Zn necessary for the present invention in the quenched ribbon sample.

Therefore, in the following examples, a soft magnetic alloy sample containing Zn in the range of 0.025 to 0.2 atomic% was produced by introducing a raw material Zn of about 20 times the desired composition into the crucible and producing a ribbon with the apparatus shown in FIG. 1. .

In addition, since it is considered that it is the same also when obtaining a thin film form, it is preferable to form into a film using the target or the evaporation source which contained much Zn which is easy to disappear beforehand.

The high saturation magnetic flux density low iron loss Fe-based soft magnetic alloy produced as described above has high saturation magnetic flux density and permeability, and the permeability is further improved according to the addition effect of Zn, and the coercivity is low, It is possible to provide a high saturation magnetic flux density low iron loss Fe-based soft magnetic alloy having a high fracture strain and bending resistance.

Next, the highly saturated magnetic flux density low iron loss Fe-based soft magnetic alloy of the present invention will be described.

One embodiment of the alloy according to the present invention mainly contains bccFe containing Fe as a main component, containing Zr, Nb and B, further containing Zn, and at least 50% or more of the structure having an average grain size of 100 nm or less. It consists of fine crystal grains, and the balance is characterized in that the amorphous alloy phase.

A more specific embodiment of the soft magnetic alloy of the present invention is composed of a composition represented by the following formula, wherein at least 50% or more of the structure is composed of microcrystal grains mainly composed of bccFe having an average grain size of 100 nm or less, and the balance is amorphous. The microcrystalline grains composed of an alloy phase, mainly composed of bccFe, are formed by quenching an alloy (melt) having a composition described below to form a nearly amorphous single phase structure, and then cooling the precipitated amorphous phase after heating above a crystallization temperature. .

(Fe _1-a Q _a ) _b Zr _x Nb _y B _t Zn _z

However, Q is either Co, Ni, or both, and a, b, x, y, t, and z, which represent the composition ratio,

0 ≤ a ≤ 0.05, 80 atomic% ≤ b, 1.5 atomic% ≤ x ≤ 2.5 atomic%, 3.5 atomic% ≤ y ≤ 5.0 atomic%, 5 atomic% ≤ t ≤ 12.5 atomic%, 0.025 atomic% ≤ z ≤ 0.200 atom %, 5.0 atomic% <x + y <7.5 atomic%, 1.5 / 6 <x / (x + y) <2.5 / 6.

Another specific aspect of the soft magnetic alloy of the present invention is composed of a composition represented by the following formula, wherein at least 50% or more of the structure consists of microcrystal grains mainly composed of bccFe having an average grain size of 100 nm or less, and the balance is amorphous. The microcrystalline grains composed of an alloy phase, mainly composed of bccFe, are formed by quenching an alloy (melt) having a composition described below to form a nearly amorphous single phase structure, and then cooling the precipitated amorphous phase after heating above a crystallization temperature. .

(Fe _1-a Q _a ) _b Zr _x Nb _y B _t Zn _z M ' _u

However, Q is either or both of Co and Ni, and M 'is one or two or more elements selected from Cr, Ru, Rh, and Ir, and represents a, b, x, y, t, z, U is

0 ≤ a ≤ 0.05, 80 atomic% ≤ b, 1.5 atomic% ≤ x ≤ 2.5 atomic%, 3.5 atomic% ≤ y ≤ 5.0 atomic%, 5 atomic% ≤ t ≤ 12.5 atomic%, 0.025 atomic% ≤ z ≤ 0.200 atom %, u ≦ 5 atomic%, 5.0 atomic% ≦ x + y ≦ 7.5 atomic%, 1.5 / 6 ≦ x / (x + y) ≦ 2.5 / 6.

The soft magnetic alloy of the above-described composition is obtained by quenching an alloy of the composition in a melt to obtain an amorphous alloy ribbon or amorphous alloy powder containing crystalline in an amorphous phase or part thereof, and heating the obtained in these processes to precipitate fine crystalline. It can be normally obtained by the heat treatment condition step.

However, in the alloy composition, as in the second soft magnetic alloy, Zn is easily lost by evaporation as compared with other elements. Therefore, the amount of Zn to be added when producing a melt as described above is set to be larger than the range of the composition formula described above. It needs to be placed. In addition, when manufacturing a thin film, it is necessary to adjust the same amount of Zn.

In addition, it is a matter of course that what was obtained by the above-mentioned quenching method may be thin film form, powder form, or thin film form, and may heat-process after shaping | molding process or machining to the shape aimed at obtained.

B is necessarily added to the soft magnetic alloy of the said composition. B has the effect of enhancing the amorphous formation method of the soft magnetic alloy, the thermal stability of the Fe-M (= Zr, Hf, Nb, etc.)-Based microcrystalline alloy, and can be a barrier to grain growth, and has a thermally stable amorphous phase. Has the effect of remaining at the grain boundaries. As a result, in the above-described heat treatment condition step, a fine body-centered cubic structure (bcc structure) having a particle diameter of 100 nm or less (specifically, 30 nm or less) that does not adversely affect magnetic properties under a wide heat treatment condition of 400 to 750 ° C. A tissue mainly composed of grains can be obtained. The content of B is preferably 5 atomic% or more and 12.5 atomic% or less, more preferably 6 atomic% or more, 9.5 atomic% or less, most preferably 8 atomic% or more and 9.0 atomic% or less. .

Further, if necessary, Y, La, Ce, Pr, Nd, Pm, Sm, Eu, Gd, Tb, Dy, Ho, Er, Tm, Yb, Lu, Cd, In, Sn, Pb, As, Sb, The magnetostriction of the soft magnetic alloy can be adjusted by adding elements such as Bi, Se, Te, Li, Be, Mg, Ca, Sr, and Ba.

In addition, since it does not interfere with the element which has amorphous forming ability, such as Al, Si, C, P, and the like, even if it contains in the range which does not reduce the characteristic of this invention alloy, content of these elements should be about 0.5 atomic% or less. desirable. Moreover, as unavoidable impurity elements, such as H, N, O, and S, to the extent that the target characteristic does not deteriorate, Preferably it does not interfere even if it contains 0.1 atomic% or less.

In the soft magnetic alloy of the above-mentioned composition, Zr and Nb preferably contain Zr and Nb having high amorphous forming ability in order to easily obtain an amorphous phase.

In addition, Zr, Nb, and the like are hardly dissolved in bccFe inherently, but the alloy melts are quenched and amorphous, and then crystallized by heat treatment conditions. The magnetostriction can be made small. In other words, the high capacity of Zr and Nb can be adjusted under heat treatment conditions, whereby the magnetostriction can be adjusted to reduce its value.

Therefore, in order to obtain low magnetostriction, it is necessary to obtain a fine crystal structure under a wide range of heat treatment conditions. As described above, it is possible to obtain a fine crystal structure under a wide range of heat treatment conditions by the addition of B. The magnetic anisotropy is brought together, resulting in good magnetic properties.

By the way, in order to reduce iron loss in particular, it is necessary to make the ratio of Zr and Nb into the prescribed range, and when Zr and Nb mentioned above are added, the relationship of 5 <(Zr content + Nb content) <7.5 is preferable, and 5.7 The range of ≤ (Zr content + Nb content) ≤ 6.5 is more preferable.

Furthermore, it is preferable to make the value of (Zr content) / (Zr content + Nb content) into the range of 1.5 / 6-2.5 / 6. This relationship is expressed by the formula: 1.5 / 6 ≤ (Zr content) / (Zr content + Nb content) ≤ 2.5 / 6, i.e. x and y representing Zr and Nb compositional ratios, 1.5 / 6 ≤ x / (x + y) ≤ 2.5 / 6.

Moreover, among the range of this relational expression, the relationship of x / (x + y) = 2/6 becomes the most preferable value.

In addition, the composition ratio x of Zr may be 0.5 atomic% or more, 3.5 atomic% or less, more preferably 1.5 atomic% or more and 2.5 atomic% or less in order to obtain good soft magnetic properties. Further, in order to obtain good soft magnetic properties, the composition ratio y of Nb is preferably 3 atomic% or more, 5.5 atomic% or less, more preferably 3.5 atomic% or more and 5.0 atomic% or less.

The corrosion resistance is improved by adding Cr, Ru, Rh, and Ir to the composition as necessary, but in order to maintain a high saturation magnetic flux density, the amount of these elements added is preferably 5 atomic% or less. In consideration of both the soft magnetic properties and the iron loss, the content of 1 atomic% or less is more preferable.

The microcrystalline structure can be obtained by partially crystallizing the Fe-M (= Zr, Hf) -based amorphous alloy by a special method as in the case of the second soft magnetic alloy described above.

Content of Fe in the soft magnetic alloy of the said composition, or each content of Fe, Co, and Ni is 80 atomic% or more, Preferably it is less than 90 atomic%. This is because a high permeability cannot be obtained when these contents exceed 90 atomic%, but in order to obtain a saturation magnetic flux density of 1.55 T or more, the range is preferably 83 to 87 atomic%, more preferably 85 to 86 atomic%. Do. Moreover, if Fe is not contained at least 80 atomic% or more, preferable saturation magnetic flux density will not be obtained.

Next, in Zn content in the soft magnetic alloy of the said composition, the range of 0.025 atomic% or more and 0.2 atomic% or less is preferable. If it is added in this range, coercive force and iron loss can be made low, and permeability can also be made high, without reducing the high saturation magnetic flux density of 1.5 T or more.

In addition, the range of 0.034 atomic% or more and 0.16 atomic% or less with respect to the content of Zn is more preferable, and in this range, a soft magnetic alloy having lower iron loss, higher saturation magnetic flux density and less change with time can be obtained. have.

However, since Zn has a melting point of 419.5 ° C. and a boiling point of 908 ° C., when melting the soft magnetic alloy of the composition described above in a crucible, the melting temperature is set to about 1240 to 1350 ° C., so most of Zn is evaporated and lost. This can be said to be the same as the second soft magnetic alloy described above.

The high saturation magnetic flux density low iron loss Fe-based soft magnetic alloy produced as described above has a high saturation magnetic flux density and permeability due to the effect of the addition of a prescribed amount of Zr and Nb, and the permeability according to the addition effect of Zn. Further, it is possible to provide a high saturation magnetic flux density, a high saturation magnetic flux density, a low iron loss Fe-based soft magnetic alloy, which is further improved and has a low coercive force and is large in fracture strain and resistant to bending.

EMBODIMENT OF THE INVENTION Below, embodiment of the low iron loss magnetic core of this invention is described in detail.

The low iron loss magnetic core according to the present invention is realized, for example, in a toroidal shape. The Trojan low iron loss magnetic core is manufactured by the quenching method of the ribbon of the low iron loss Fe-based soft magnetic alloy of the composition described above, press-punched the ribbon to obtain a ring, and the ring is laminated in the required number of sheets. Or a ribbon of low iron loss Fe-based soft magnetic alloy to form an annular shape, and the core is coated with a resin of epoxy resin or encapsulated in a resin case for insulation protection. Obtained.

In order to realize the core of the E1 core type, the ribbon is press-punched to form the E type or the I type, and a plurality of the E type thin films and the I type thin films are formed, and then the E type thin films or the I type thin films are formed. It can obtain by forming an E type core and an I type core by laminating | stacking and joining them. In addition, the E-type cores and I-type cores are epoxy-based resins, for example, resin-coated or inserted into a resin case to insulate and protect the necessary portions, and after winding, the sides of the E-type magnetic core and type I The low iron loss magnetic core is obtained by joining the side portions of the magnetic core of. The magnetic core is not limited to a combination of an E type and an I type core, and may be a magnetic core formed of any combination of an E type core and an E type core, a U type core and an I type core, and a U type and a U type core. No problem

3 and 4 show an example of a trolley-shaped low iron loss magnetic core. In the configuration shown in FIG. 3, the present invention is provided inside the upper and lower case 21 and the hollow case of the hollow annular shape. A low iron loss magnetic core in which a ring of a low iron loss Fe-based soft magnetic alloy ribbon related to each other is laminated is stored. In the configuration shown in FIG. 4, the low iron loss Fe is formed inside the same upper case 20 and lower case 21. A magnetic core is formed by storing the low iron loss magnetic core 24 having the resin-coated whole wound around the ribbon 23 of the magnetic magnetic alloy. In addition, the upper case 20 and the lower case 21 may be used suitably, and of course, you may comprise a magnetic core only by resin coating.

5 shows an example in which the low iron loss magnetic core of the present invention is applied to a common mode choke coil.

The common mode choke coil 25 is used for noise filter or the like as three commercially available magnetic cores 26, which is formed by winding a ribbon of a low iron loss Fe-based soft magnetic alloy according to the present invention and coating the resin with the magnetic core 26. It consists of three windings 27 wound around each and a bobbin 28 mounted on a magnetic core 26.

The low iron loss Fe-based soft magnetic alloy constituting the ribbon has a bcc-Fe phase (body core) of 100 nm or less, more preferably 30 nm or less, having an average grain size of 50% or more, preferably 70% or more of the structure. It is composed mainly of a microcrystalline structure mainly composed of microcrystalline grains mainly composed of a cubic Fe phase) and a residual amorphous structure, and thus exhibits little magnetostriction, high saturation magnetic flux density and excellent transmittance.

The low iron loss Fe-based soft magnetic alloy according to the present invention is a step obtained by rapidly quenching an alloy mainly composed of the amorphous phase of the above composition in a melt, a microcrystalline structure composed of fine crystal grains by heating and cooling the alloy obtained in this step. It can obtain normally by the heat processing conditions which precipitate a. The composition of these low iron loss Fe-based soft magnetic alloys can be preferably applied by applying the above-described first to third low iron loss Fe-based soft magnetic alloys.

Since the low iron loss magnetic core made of such a soft magnetic alloy has excellent soft magnetic properties and low iron loss, it can be used as a magnetic core of various transformers such as low frequency transformers, pulse transformers, power transformers, columnar transformers, etc. It can be used suitably as a magnetic core of a coil.

In the low iron loss magnetic core described above, at least 50% of the structure is a microcrystalline structure mainly composed of a bcc-Fe phase having an average grain size of 100 nm or less, and preferably the first to third low iron loss Fe bases. Since the soft magnetic alloy has high transmittance, high saturation magnetic flux density, and low iron loss, magnetic loss of low iron loss made of such a soft magnetic alloy is used in a transformer, for example, to reduce power energy loss. At the same time, it is possible to reduce the amount of heat generated by the magnetic core itself.

Furthermore, since the fracture strain of the low iron loss Fe-based soft magnetic alloy according to the present invention is larger than 1.0 × 10 ⁻² or more, for example, when the shape of the soft magnetic alloy is used as a ribbon, the ribbon has excellent bending workability. As a result, the ribbon can be easily wound to form a ring-shaped magnetic core.

Next, the first manufacturing method of the present invention will be described.

In the present invention, the amorphous alloy ribbon produced by the manufacturing apparatus shown in Fig. 1 is heat-treated to precipitate a microcrystalline structure mainly composed of bccFe. In the first manufacturing method of the present invention, the temperature increase rate until reaching a predetermined heat treatment condition temperature is 10 ° C / min or more and 200 ° C / min or less, more preferably 20 ° C / min or more and 100 ° C / min or less. Preferably it is 20 degreeC / min or more and 40 degrees C / min or less. In addition, when the temperature increase rate is low, there is a problem in that the process water is large or sufficient magnetic properties are not obtained. On the contrary, when the temperature is too fast, the heat treatment furnace cannot cope with a high temperature rise temperature or when the heat treated alloy becomes large, This becomes difficult to transfer, which makes the conduction of heat uneven, making it difficult to obtain a uniform crystal structure, leading to deterioration of magnetic properties, which is undesirable. In addition, in order to heat-treat a large alloy at a uniform and high temperature increase rate, a complicated structure is required in the heat treatment furnace, resulting in high cost of a manufacturing facility.

In the present invention, it is possible to find a temperature increase rate that can obtain a good magnetic property and to obtain a critically excellent soft magnetic property, even if the temperature increase rate is not particularly increased, and to apply it to the manufacturing method of the present invention.

By the way, with respect to the heat treatment temperature, in accordance with the following regulation generally in the present invention, it is possible to produce an alloy of high properties. That is, when the crystallization temperature of bccFe is T _X1 [° C] and the crystallization temperature of the compound phase crystallized on the higher temperature side than the T _X1 is T _X2 [° C], the temperature interval T _x of both of these T _X1 and T _X2 is T _X2 -T _X1 ) is preferably made as large as possible. Specifically, it is appropriate that a T _x ≥200 ℃. The reason for this is based on the viewpoint of obtaining a microcrystalline grain of bccFe mainly composed of an amorphous alloy phase as an alloy in the present invention, and avoiding unnecessary crystallization of the crystal phase by performing heat treatment up to T _X2 . Because you have to. In this case, the interval between the crystallization temperature between the bccFe phase and the compound phase is widened. Therefore, only the bccFe phase is precipitated by heat treatment of the alloy under optimum conditions to suppress the precipitation of other compound phases. It is possible to improve the soft magnetic properties of the. Specifically, when the above-mentioned heat treatment temperature is set to Ta, it is preferable to satisfy the condition of T _X1 <Ta <T _X2 . In addition, here, the compound phase crystallized in said T _X2 is the high saturation magnetic flux density low iron loss Fe base soft magnetic alloy in the case of the 1st-3rd case demonstrated below specifically, Fe ₃ B, Fe ₂ It is considered to be B.

In addition, in the intermediate value T _X1 [° C] between the T _X1 and T _X2 , crystallization of a separate compound phase is sometimes performed. In the present invention, the composition of the separate phase is not clear, but the precipitation of the separate phase tends to be influenced by the composition of the highly saturated magnetic flux density low iron loss Fe-based soft magnetic alloy, and especially when the composition ratio of B is high, Since it is easy, it is considered to be a compound phase of B and other additive elements. However, this is not clear. If this observation is the crystallization temperature (T _X1) of the "phase separate", the as T _X, this T _X1 - _X1 by T, and T is _X is not less than 200 ℃, it is preferred that. The reason is as described above. In addition, the heat treatment temperature Ta of the alloy must also satisfy T _X1 <Ta <T _X1 .

The high saturation magnetic flux density low iron loss Fe-based soft magnetic alloy in the present invention is produced by the above manufacturing apparatus, but the composition of the same alloy that can be applied to the first manufacturing method is the first to third described above. It is preferable to apply to the highly saturated magnetic flux density low iron loss Fe-based soft magnetic alloy of.

Next, the second manufacturing method of the present invention will be described.

The single-roll liquid quenching method using the apparatus shown in FIG. 1 can be used preferably. That is, molten metal or melt of a predetermined component is ejected by a pressure of Ar gas and quenched from a quartz nozzle placed on one rotating steel cooling roll to obtain a ribbon. At this time, the temperature of the melt ejected from the nozzle, that is, the injection temperature, is less than 1350 ° C in this embodiment. However, as a comparative example, the alloy was also manufactured in the case where the injection temperature was at least 1350 ° C. If the injection temperature is too low, the viscosity of the melt may be lowered and the nozzle may be clogged. Therefore, in the present invention, the temperature should be set to 1240 ° C or higher. It should be noted that the above-mentioned "molten metal of the predetermined component" corresponds to the highly saturated magnetic flux density low iron loss Fe base magnetic alloys each having the first to third compositions described above. However, the following attention is required about the 2nd and 3rd case containing Zn. In other words, when the melting point and boiling point of Zn are 419.5 ° C and 908 ° C, respectively, and the injection temperature is around 1350 ° C, more specifically, about 1240-1350 ° C, most of Zn is evaporated and disappeared. Therefore, in order to finally secure the composition ratio of Zn as an alloy that satisfies the composition ratio in the second and third cases, it is necessary to inject Zn in an amount exceeding the above Zn amount into the crucible. Specifically, in order to make the target composition, that is, 0.025 atomic% or more and 0.2 atomic% or less, the raw material Zn of about 20 times may be added to the crucible.

Next, the third manufacturing method of the present invention will be described.

In producing the Fe-based soft magnetic alloy by the third production method of the present invention, Fe is first selected from the group consisting of Ti, Zr, Hf, V, Nb, Ta, Mo, W, and Mn as main components. Alternatively, the alloy melt containing the elements M and B consisting of two or more metal elements is quenched to produce an amorphous alloy ribbon. As the manufacturing method of this alloy ribbon, a well-known method, such as injecting an alloy melt into a moving cooling body, such as a cooling roll rotating at high speed, using the manufacturing apparatus as shown in FIG. 1, for example can be used.

Subsequently, the produced amorphous alloy ribbon is subjected to a first heat treatment at a holding temperature of 500 to 800 ° C.

The quenched alloy thin film is composed mainly of an amorphous structure, and when heated and heated, a microcrystalline phase mainly composed of Fe grains having a fine bcc structure (body-centered cubic structure) having an average grain size of 30 nm or less at a certain temperature or more. Will precipitate. In this specification, the temperature which the microcrystalline phase of Fe which has this bcc structure precipitates is called 1st crystallization temperature. This first crystallization temperature changes depending on the composition of the alloy, which is about 480 to 550 ° C.

When Fe ₃ B or Zr is contained in a temperature higher than the first crystallization temperature, a compound phase (second crystal phase) that deteriorates soft magnetic properties such as Fe ₃ Zr is precipitated. In this specification, the temperature at which such a compound phase precipitates is called 2nd crystallization temperature. This second crystallization temperature changes depending on the composition of the alloy, which is about 740 to 810 ° C.

Therefore, in the present invention, the holding temperature at the time of performing the first heat treatment on the amorphous alloy ribbon is preferably in the range of 500 ° C to 800 ° C, whereby the microcrystalline phase containing Fe as the main component having a bcc structure is preferably precipitated. In order not to precipitate, it sets suitably according to the composition of an alloy.

In the present invention, the time for maintaining the amorphous alloy ribbon at the holding temperature can be a short time of 20 minutes or less, and depending on the composition of the alloy, 0 minutes, i.e., the temperature is immediately lowered after raising the temperature, even if there is no holding time, high permeability Can be obtained. Moreover, especially in the case of a composition which does not contain Cu and Si, especially Si, high permeability can be obtained with a shorter holding time of 10 minutes or less. This is because, when Si is added, it is necessary to sufficiently dissolve Si in Fe, so the holding time must be extended. Although the holding time may be longer than the above range, the magnetic properties are not improved even if the holding time is longer, and the production time is longer, which is not preferable.

Moreover, the temperature increase rate which raises the temperature of an amorphous alloy ribbon from room temperature to the said holding temperature at the time of 1st heat processing is 10 degreeC or more, More preferably, it is 10-200 degreeC / min or less, More preferably, 30 degreeC / It becomes more than minutes and 100 degrees C / min or less. The slower the temperature increase rate, the longer the production time is, so the higher the temperature increase rate is desirable, it is difficult to increase the size of the current heating apparatus more than about 200 ℃ / min.

After the first heat treatment is performed, the alloy ribbon is cooled to a predetermined temperature by air cooling or the like, and the second heat treatment is performed, and then the alloy ribbon is cooled to room temperature by air cooling or the like. In the present invention, the method of performing the second heat treatment during the cooling after the first heat treatment is called two-stage annealing.

The predetermined temperature here refers to the holding temperature at the time of performing the second heat treatment. The holding temperature of the second heat treatment is 100 ° C or higher and a temperature below the holding temperature of the first heat treatment, preferably 200 to 400 ° C. If the holding temperature of the second heat treatment is less than 100 ° C, the annealing effect is hardly obtained, and thus the soft magnetic properties cannot be sufficiently improved. In addition, if the holding temperature of the second heat treatment exceeds the holding temperature of the first heat treatment, when Fe ₃ B or Zr is contained in the alloy, the compound phase deteriorates soft magnetic properties such as Fe ₃ Zr. 2 crystal phase) is precipitated.

In the second heat treatment, the time for holding the alloy ribbon at the holding temperature is 0.5 to 100 hours, preferably 1 to 30 hours. If the time (holding time) to be maintained at the holding temperature at the time of performing the second heat treatment is less than 0.5 hour, the coercive force is so large that soft magnetic properties such as permeability cannot be sufficiently improved. The change over time becomes larger.

In this way, by cooling the heat treatment in two stages, the soft magnetic properties can be further improved, and a Fe-based soft magnetic alloy having a small change with the passage of the magnetic properties over time even in a high temperature state is obtained.

In addition, the first heat treatment may be performed by lowering the alloy ribbon to room temperature by air cooling or the like, and then raising the temperature at the predetermined temperature to perform the second heat treatment. Here, the second heat treatment that is heat treated at a temperature lower than the first heat treatment is referred to as low temperature annealing. Thus, by depositing the crystal phase which consists of crystal grains of the fine bcc structure whose average grain diameter is 30 nm or less in an amorphous alloy by a 1st heat treatment, by carrying out low temperature annealing, soft magnetic characteristics can be improved further and it is high temperature state. Even if left for a long time, a Fe-based magnetic alloy having a small change in magnetic properties over time is obtained.

Examples of the heat treatment pattern in the method for producing the Fe-based soft magnetic alloy of the present invention include those shown below.

As a heat treatment pattern in the case of performing two-stage annealing, for example, as shown in the graph of FIG. 6, after the amorphous alloy ribbon is heated from room temperature to the holding temperature of the first heat treatment, the holding at the constant temperature within the holding temperature range. After holding for some time, the temperature is lowered to the holding temperature of the second heat treatment, and after holding the holding time at a constant temperature within the holding temperature range, a pattern for lowering the alloy ribbon to room temperature by air cooling or the like is mentioned.

As a heat treatment pattern in the case of performing the second heat treatment followed by the second heat treatment and low temperature annealing, for example, as shown in the graph of FIG. 7, after raising the amorphous alloy ribbon from room temperature to the holding temperature of the first heat treatment, After the above holding time is maintained at a constant temperature within the holding temperature range, the alloy ribbon is lowered to room temperature by air cooling or the like, and then the alloy ribbon is raised from room temperature to the holding temperature of the second heat treatment, and the constant temperature within the holding temperature range is maintained. After maintaining the said holding time at temperature, the pattern which cools the said alloy ribbon to room temperature by air cooling etc. is mentioned.

In the method for producing the Fe-based soft magnetic alloy of the present invention, by performing the first heat treatment on the amorphous alloy ribbon as described above, the average crystal grain size is 30 nm without precipitating a compound phase that deteriorates soft magnetic properties such as Fe ₃ B. A microcrystalline alloy containing an amorphous phase is obtained mainly based on the crystal grains of Fe having the following fine bcc structure. Thus, excellent soft magnetic characteristics can be exhibited by using a structure mainly composed of a crystal phase composed of fine crystal grains and a grain boundary amorphous phase present at the grain boundary.

Further, the microcrystalline alloy is subjected to a second heat treatment at a holding temperature of 100 ° C. or higher and a holding temperature of the first heat treatment temperature or lower, thereby providing better soft magnetic properties, and even when left in a high temperature state for a long time. A Fe-based magnetic alloy with less change over time was obtained.

As the reason why the alloy produced according to the present invention exhibits excellent soft magnetic properties, the crystals which were regarded as one of the causes of deterioration of the soft magnetic properties in the conventional crystalline material because of the small particle diameter of the bcc crystal grains deposited by the first heat treatment. It is considered that the magnetic anisotropy is averaged by the magnetic interaction between the bcc particles, and the crystallite anisotropy in appearance is very small.

Here, if the average grain size of the crystal grains to be the main body is larger than 30 nm, it is not preferable because averaging of the crystal magnetic anisotropy is insufficient and the soft magnetic characteristics deteriorate.

In addition, it is considered that the residual stress in the sample generated in the first heat treatment is alleviated by the second heat treatment.

The Fe-based soft magnetic alloy is preferably one containing Fe as the main component, and elements M and B composed of one or two or more selected from Ti, Zr, Hf, V, Nb, Ta, Mo, W, and Mn. . Specifically, the third manufacturing method of the present invention can be suitably applied to the first to third highly saturated magnetic flux density low iron loss Fe-based soft magnetic alloys described above.

Or it is suitable also for the soft magnetic alloy of the composition shown by each following formula.

(Fe _1-a Z _a ) _b B _x M _y

(Fe _1-a Z _a ) _b B _x M _y X _z

(Fe _1-a Z _a ) _b B _x M _y T _t

(Fe _1-a Z _a ) _b B _x M _y T _t X _z

Provided that Z is one or two or more elements of Ni and Co, M is one or two or more elements selected from Ti, Zr, Hf, V, Nb, Ta, Mo, W, and Mn, and T is Cu, One or two or more elements selected from the group consisting of Ag, Au, Pd, and Pt, X is one or two or more of Si, Al, Ge, Ga, and includes a, b, x, y, t, z Is 0 ≤ a ≤ 0.1, 75 atomic% ≤ b ≤ 93 atomic%, 0.5 atomic% ≤ x ≤ 18 atomic%, 4 atomic% ≤ y ≤ 9 atomic%, t ≤ 5 atomic%, z ≤ 5 atomic%. It goes without saying that these can also be preferably applied to the first and second production methods as well.

In addition, the above-mentioned 1st thru | or 3rd manufacturing method may apply either one, or may combine and apply two or more.

Example

Example 1 Structure of Highly Saturated Magnetic Flux Density Low Iron Loss Fe-based Soft Magnetic Alloy

First, an alloy ribbon mainly containing an amorphous phase is prepared by the single-roll liquid quenching method. That is, the melt is ejected onto the roll by the pressure of argon gas from the nozzle placed on one rotating steel roll and quenched to obtain a ribbon. The width of the ribbon created as mentioned above is about 15 mm, and the thickness is about 20 micrometers.

Next, the obtained ribbon was heat-treated under vacuum at a heating rate of 180 ° C./min, a heat treatment temperature of 535 ° C., and a holding time of 5 minutes to precipitate microcrystalline structures to form a ribbon of highly saturated magnetic flux density low iron loss Fe-based magnetic alloy.

With respect to the ribbon of the obtained soft magnetic alloy, the state of the structure of the ribbon is examined by X-ray diffraction measurement. In addition, the magnetic permeability (μ '), coercive force (Hc), and saturation magnetic flux density (B ₁₀ ) of the ribbon of the soft magnetic alloy were measured.

Permeability is measured by using an inductance analyzer by processing a ribbon into a ring shape having an outer diameter of 10 mm and an inner diameter of 6 mm, winding the stacked one. The measurement conditions of the magnetic permeability (μ ') are 5 m Oe and 1 kH _Z. Coercive force (Hc) and saturation magnetic flux density (B ₁₀ ) are measured using a direct current BH loop tracer.

First, the effect of the heat treatment on the structure of the highly saturated magnetic flux density low iron loss Fe-based soft magnetic alloy of the present invention will be described by taking a ribbon of a soft magnetic alloy having a composition of Fe _85.75 Zr ₂ Nb ₄ B _8.25 .

The change of the structure before and after the heat treatment of the ribbon of the soft magnetic alloy of the composition is investigated by X-ray diffraction method. The results are shown in FIGS. 8 and 9.

In Fig. 8, in the quenched state (the state in which the melt is quenched and turned into a ribbon), the halo diffraction pattern peculiar to amorphous is shown. In Fig. 9, the diffraction pattern peculiar to Fe of the body centered cubic crystal (bcc) after heat treatment, It was confirmed that the structure of the alloy changed from amorphous to body-centered cubic crystal by heat treatment. In addition, the present inventors observed the state of the structure using a transmission electron microscope, and confirmed that the structure after heat treatment consists of microcrystals having a particle size of about 10 nm.

The alloy of the present invention as described above can be crystallized by heat treatment of the amorphous alloy having the above-described composition to obtain a microcrystalline structure mainly containing ultrafine grains.

Fe ₈₆ Nb ₇ B ₇ , Fe ₉₁ Zr ₇ B ₂ , Fe ₈₉ Zr ₇ B ₄ , Fe ₈₉ Zr ₅ B ₆ (above, comparative example), Fe ₈₆ Zr ₂ Nb ₄ B ₈ , Fe _85.75 Zr ₂ Nb ₄ B _8.25 , Fe _85.5 Zr ₂ Nb ₄ B _8.5 to prepare a quench ribbon, the temperature is raised at a heating rate of 180 ℃ / min, subjected to a heat treatment for 5 minutes to 1 hour at a temperature of 510 ~ 650 ℃ lead Obtain a magnetic alloy. The magnetic permeability (μ '), coercive force (Hc) and saturation magnetic flux density (B ₁₀ ) of these soft magnetic alloys are measured. The results are shown in Table 1.

As shown in Table 1, the composition of Fe ₈₆ Zr ₂ Nb ₄ B ₈ , Fe _85.75 Zr ₂ Nb ₄ B _8.25 , Fe _85.5 Zr ₂ Nb ₄ B _{8.5 according to} the present invention was added at the same time by adjusting the composition ratio of Zr and Nb. One soft magnetic alloy has a high permeability (μ '), a low coercive force (Hc), and excellent soft magnetic properties, compared to an alloy system in which Zr and Nb are added alone. In particular, an alloy having a composition of Fe _85.75 Zr ₂ Nb ₄ B _8.25 has a magnetic permeability (μ ') of 57800 and a coercive force (Hc) of 0.043 Oe, showing particularly excellent soft magnetic properties.

Example 2 Relationship between Alloy Composition and Various Properties

Next, in the same manner as in Example 1, an alloy ribbon in a quenched state was formed and heat-treated again to prepare a soft magnetic alloy having various compositions. In addition, as long as there is no restriction | limiting in particular, the ribbon heat processing was performed on the conditions of the temperature increase rate of 180 degree-C / min, holding time 5 minutes.

For the ribbon of the obtained soft magnetic alloy, the coercive force (Hc), the magnetic permeability (μ ') at 1 kHz, the saturation magnetic flux density (B ₁₀ ) and the residual magnetization (Br) in the magnetic field of 10 Oe are measured. In addition, the constant (λ _s ) of the magnetostriction is measured for a part of the obtained ribbon.

Moreover, when heat-processing in 500 degreeC-700 degreeC, the heat processing temperature which the coercive force (Hc) becomes the minimum and the magnetic permeability (micro ') becomes the maximum is measured.

In addition, about a part of obtained ribbon, the average grain size of the crystal grain in a microcrystal structure is calculated | required by X-ray diffraction method.

In addition, differential thermal analysis (DTA measurement) was performed on a part of the ribbon in the quenched state before the heat treatment to determine the crystallization temperatures (T _X1 , T _X2 , T _{X1 ′} ) of the bcc-Fe phase and other compound phases. By measuring, the interval ΔT _X of the crystallization temperature is obtained.

These results are shown in FIGS. 10-59.

10 to 13 show various characteristics of the soft magnetic alloy in which the total amount of Zr and Nb is 5 atomic%, and FIGS. 14 to 23 show the soft magnetic alloy in which the total amount of Zr and Nb is 5.5 atomic%. 24 to 35 show various characteristics of the soft magnetic alloy in which the total amount of Zr and Nb is 6 atomic%, and FIGS. 36 to 47 show the soft magnetic alloy in which the total amount of Zr and Nb is 6.5 atomic%. 48 to 59 show characteristics of the soft magnetic alloy in which the total amount of Zr and Nb is 7 atomic%.

10 to 59, The mark shows the ribbon in which the diffraction peak of the (200) plane on the bcc-Fe was confirmed in the ribbon in the quenched state, The marks indicate ribbons in which the diffraction peaks of the (200) plane on the bcc-Fe phase were not found in the ribbon in the quenched state.

In other words, The ribbon of the indication is heat-treated the quenching ribbon in which the crystalline phase precipitated in a part of the amorphous phase, The ribbon of mark is heat-processed the quenching ribbon of substantially amorphous single phase.

`` Zr + Nb = 5 atomic% soft magnetic alloy ''

As shown in FIG. 10, the soft magnetic alloy having Zr + Nb = 5 atomic% exhibits a coercive force (Hc) of 59 to 1055 m Oe.

Here, the coercive force (Hc) of 59 m Oe is an alloy having a composition of Fe ₈₇ Zr _2.5 Nb _2.5 B ₈ .

Next, as shown in FIG. 11, the soft magnetic alloy whose Zr + Nb = 5 atomic% has shown the permeability (micro ') of 300-33000.

Here, the alloy having a composition of Fe ₈₇ Zr _2.5 Nb _2.5 B ₈ is shown to have a permeability (μ ') of 33000.

Next, there is shown an, Zr + Nb = 5 atomic percent, soft magnetic alloy is 1.59 to 1.72 T in saturation magnetic flux density (B ₁₀₎ as shown in Fig. That is, Zr is 1 atomic% or more and 2.5 atomic% or less, B is 6.75 atomic% or more and 11 atomic% or less, and the sum of Fe and Nb is 88 atomic% or more and 90.75 atomic% or less (Fe is 84 atomic% or more and 88.5 atomic% or less) In the phosphorus composition range, the saturation magnetic flux density (B ₁₀ ) of 1.5T or more is shown. Moreover, the alloy of the composition of Fe ₈₇ Zr _2.5 Nb _2.5 B ₈ has shown the high value of 1.72T.

As shown in FIG. 13, the soft magnetic alloy having Zr + Nb = 5 atomic% shows residual magnetization (Br) of 0.47 to 1.36 T. As shown in FIG.

As described above with reference to FIGS. 10 to 13, when the total amount of Zr and Nb is 5 atomic%, Zr is 1 atomic% or more and 2.5 atomic% or less, B is 6.75 atomic% or more and 11 atomic% or less, Fe and Nb It was found that the soft magnetic properties were excellent when the sum of was at least 88 atomic% and at most 90.75 atomic%. The most preferred composition was an alloy having a composition of Fe ₈₇ Zr _2.5 Nb _2.5 B ₈ .

`` Zr + Nb = 5.5 atomic percent soft magnetic alloy ''

As shown in FIG. 14, the soft magnetic alloy having Zr + Nb = 5.5 atomic% exhibits a coercive force (Hc) of 94 to 211 m Oe.

Here, the coercive force (Hc) of 200 mOe or less is an alloy of a composition range in which Zr is 1 atomic% or more and the total of Fe and Nb is 90 atomic% or less.

Moreover, Zr is preferably 1.5 atomic% or more, and the alloy of the composition range whose total of Fe and Nb is 88.5 atomic% or less is shown to show the coercive force (Hc) of 100 mOe or less.

Next, as shown in FIG. 15, the soft magnetic alloy having Zr + Nb = 5.5 atomic% has a magnetic permeability (μ ') of 8400 in 25400.

It can be seen that the magnetic permeability μ 'depends on the composition ratio of Fe, Zr and Nb, and does not depend on the composition ratio of B. Specifically, the total amount of Fe and Nb is 90.5 atomic% or less, Zr is 0.5 atomic% or more, and a permeability (μ ') of 10000 or more is obtained, the total amount of Fe and Nb is 89 atomic% or less, and Zr is 1 atomic% or more It can be seen that the permeability (μ ') of more than 20000 is obtained.

The soft magnetic alloys shown in Figs. 14 and 15 both have high magnetic permeability (μ ') and low coercive force (Hc) and exhibit excellent soft magnetic properties.

Next, as illustrated in FIG. 16, the soft magnetic alloy having Zr + Nb = 5.5 atomic% exhibits a saturation magnetic flux density (B ₁₀ ) of 1.60 to 1.68 T. Next, as shown in FIG.

In addition, as shown in FIG. 17, the soft magnetic alloy having Zr + Nb = 5.5 atomic% exhibits a residual magnetization (Br) of 0.44 to 0.62T.

FIG. 18 shows the optimum heat treatment temperature for obtaining the minimum coercive force Hc in the soft magnetic alloy having Zr + Nb = 5.5 atomic%.

Among these, the coercive force (Hc) decreases at a heat treatment temperature of 550 ° C. or lower is a ribbon in which Zr is 1 atomic% or more, B is 10 atomic%, and the sum of Fe and Nb is 88.5 atomic% or more and 89 atomic% or less. Able to know.

In addition, the heat treatment temperature is advantageously low temperature for mass production, the soft magnetic alloy of the present invention is 550 ℃ or less, it can be suppressed to a lower temperature than the conventional microcrystalline alloy described above.

19, the optimum heat treatment temperature for obtaining the maximum permeability (micro ') is shown in the soft magnetic alloy which is Zr + Nb = 5.5 atomic%.

Among them, the permeability (μ ') is increased at the heat treatment temperature of 550 ° C. or lower, which is a ribbon in which Zr is 1 atomic% or more, B is 10 atomic%, and the sum of Fe and Nb is 88.5 atomic% or more and 89 atomic% or less. Able to know.

18 and 19, the ribbon of the soft magnetic alloy having Zr of 1 atomic% or more, B of 10 atomic% and a total of Fe and Nb of 88.5 atomic% or more and 89 atomic% or less has a heat treatment temperature of 550 ° C. Even below, it can be seen that the optimum coercive force (Hc) and permeability (μ ') are exhibited, and that the soft magnetic alloy has fine crystals with excellent productivity. This composition range includes approximately the optimum composition range of the coercive force Hc and the magnetic permeability μ 'shown in FIGS. 14 and 15.

Next, Figure 20 shows the crystallization temperature (T _X1) on the bcc-Fe, shows the crystallization temperature (T _X1) on the other compounds in Figure 21, Figure 22 also shows the crystallization temperature (T _X2) on the other compound. The relationship between these crystallization temperatures is T _X1 <T _{X1 '} <T _X2 . Further, Fig. 16, the distance between the crystallization temperature: - represents the _{(ΔT X ΔT X = T X2} T X1).

As shown in FIG. 20, T _X1 is in the range of 462 to 484 ° C., and depends on the composition ratio of Fe, Nb, and Zr, and hardly depends on the composition ratio of B. As shown in FIG.

Further, the interval (ΔT _X) of the crystallization temperature as shown in Fig. 23, it indicates the range of 331 to 337 ℃. As such, these alloys exhibit a wide crystallization temperature interval (ΔT _X ) of 320 ° C. or higher, so that only the bcc-Fe phase is precipitated during heat treatment to suppress the precipitation of the compound phase, and the soft magnetic alloy deteriorates in soft magnetic properties. Can be prevented.

21, the plot without subscripts shows an alloy in which no crystallization temperature (T _{X1 '} ) of other compound phases was observed, and it can be seen that an alloy in which T _X1' does not exist usually has good magnetic properties. .

As described above with reference to FIGS. 14 to 23, when the total amount of Zr and Nb is 5.5 atomic%, Zr is 0.5 atomic% or more, preferably 1 atomic% or more, and B is 10 atomic%, When the total amount of Fe and Nb is 88.5 atomic% or more and 89 atomic% or less (Fe is 84.5 atomic% or more and 87.75 atomic% or less), excellent soft magnetic properties are exhibited.

Specifically, the highly saturated magnetic flux density low iron loss Fe-based soft magnetic alloy having a composition of Fe _84.5 Zr ₁ Nb _4.5 B ₁₀ and Fe _84.5 Zr _1.5 Nb ₄ B ₁₀ exhibits excellent soft magnetic properties.

`` Zr + Nb = 6 atomic percent soft magnetic alloy ''

As shown in FIG. 24, the soft magnetic alloy whose Zr + Nb = 6 atomic% has shown the coercive force (Hc) of 38-8400 mOe.

Here, the coercive force (Hc) of 70 mOe or less indicates that Zr is 0.5 atomic% or more, preferably 1 atomic% or more, B is 10 atomic% or less, and the composition of Fe and Nb is 90 atomic% or less. Range of alloys.

Moreover, Zr is 1.5 atomic% or more and 3.5 atomic% or less, B represents 6.5 atomic% or more and 9.5 atomic% or less, Preferably it is 6.5 atomic% or more and 9 atomic% or less to show the coercive force (Hc) of 50 mOe or less. And alloy of composition range whose total of Fe and Nb is 89 atomic% or more and 90 atomic% or less (Fe is 84.5 atomic% or more and 87.5 atomic% or less).

In addition, the coercive force (Hc) of 40 mOe or less indicates that Zr is 1.5 atom% or more and 2.5 atom% or less, preferably 2.0 atom%, and B is 8 atom% or more and 9 atom% or less, and Fe and Nb It is an alloy of the composition range whose sum total is 89 atomic% or more and 89.5 atomic% or less, Preferably it is 89.5 atomic% (Fe is 85 atomic% or more and 86 atomic% or less).

Next, as shown in FIG. 25, the soft magnetic alloy whose Zr + Nb = 6 atomic% has shown the magnetic permeability (micro ') of 900-59000.

The permeability (μ ') of 30000 or more here is an alloy of the composition range whose Zr is 1 atomic% or more, B is 10 atomic% or less, and the sum total of Fe and Nb is 90 atomic% or less.

Moreover, Zr is 1 atomic% or more and 3 atomic% or less, B represents 7.5 atomic% or more and 9.5 atomic% or less, and the sum total of Fe and Nb is 89 atomic% or more and 90 atomic% (Fe is an alloy of the composition range of 84.5 atomic% or more and 86.5 atomic% or less).

Moreover, Zr is 1.5 atomic% or more and 2.5 atomic% or less, B is 8 atomic% or more and 9 atomic% or less, and the sum total of Fe and Nb is 89 atomic% or more 90 which shows the magnetic permeability (micro ') of 50000 or more. It is an alloy of the composition range which is atomic% or less (Fe is 85 atomic% or more and 86 atomic% or less).

Therefore, as can be seen from Figs. 24 and 25, Zr + Nb = 6 atomic% is satisfied, and Zr is 1.5 atomic% or more and 2.5 atomic% or less, that is, the range of Zr / (Zr + Nb) is 1.5 / 6 or more and 2.5 / 6 or less, B is 8 atomic% or more and 9 atomic% or less, Fe is 80 atomic% or more, and the sum of Fe and Nb is 89 atomic% or more and 90 atomic% or less (Fe is 85 atoms The ribbon of the soft magnetic alloy (% or more and 86 atomic% or less) has a high magnetic permeability (μ ') of 40000 to 50000 or more and a low coercive force (Hc) of 40 mOe or less, and shows excellent soft magnetic properties.

Next, as shown in FIG. 26, the soft magnetic alloy Zr + Nb = 6 atomic percent, shows the saturation magnetic flux density (B ₁₀₎ of 1.53 to 1.67 T. The saturation magnetic flux density (B ₁₀ ) depends on the composition ratio of Fe, and the correlation with the composition ratio of Zr, Nb, and B is not clear. However, when the concentration of Fe is high, the saturation magnetic flux density (B ₁₀ ) also tends to increase. and, if it is possible to have with the above-mentioned composition range, 1.5 to 1.6 T or more saturation magnetic flux density (B ₁₀₎ and 40000 to the high magnetic permeability (μ ') at least 50,000.

As shown in FIG. 27, the soft magnetic alloy having Zr + Nb = 6 atomic% exhibits residual magnetization (Br) of 0.39 to 1.19 T. As shown in FIG. Regarding the residual magnetization (Br), the correlation with the composition ratio of Fe, Zr, Nb, and B is not clear.

As shown in Fig. 28, in the soft magnetic alloy having Zr + Nb = 6 atomic%, it can be seen that the average grain size of the bcc-Fe phase is very fine, 10 to 12 nm.

Among them, the average grain size is 11 nm or less, Zr is 4 atomic% or less, B is 5.5 atomic% or more and 10 atomic% or less, preferably 6 atomic% or more and 9 atomic% or less, and the sum of Fe and Nb. Is 88 atomic% or more and 92 atomic% or less, preferably 89 atomic% or more and 92 atomic% or less (Fe is 84 atomic% or more and 88.5 atomic% or less, preferably 85 atomic% or more and 88 atomic% or less). It is a ribbon, and it turns out that these ribbons have a especially fine crystalline structure. In addition, this composition range includes the optimum composition range of the coercive force (Hc) and the permeability (μ ') shown in FIGS. 24 and 25. It is suggested that the finer the average grain size of the bcc-Fe phase, the better the soft magnetic characteristics. do.

As shown in Fig. 29, in the soft magnetic alloy having Zr + Nb = 6 atomic%, the magnetostriction has a constant (λ _s ) in the range of -14 x 10 ^-7 to 17 x 10 ^-7 , and a good value. Indicates. In addition, the luxury line which has a magnetostriction of 0 is contained in the region with the highest permeability (micro ') shown in FIG. The constant (λ _s ) of the magnetostriction tends to depend on the composition ratio of B. The constant B of the magnetostriction (λ _s ) is almost zero when B is 8 or more and 9 or less.

Fig. 30 shows the optimum heat treatment temperature for obtaining the minimum coercive force Hc in the soft magnetic alloy having Zr + Nb = 6 atomic%.

In particular, the coercive force (Hc) decreases at a heat treatment temperature of 525 ° C. or lower, where Zr is 1 atomic% or more and 3 atomic% or less, and B is 7.5 atomic% or more and 9.5 atomic% or less, and Fe and Nb It turns out that the sum total is 89 atomic% or more and 90 atomic% or less (Fe is 84.5 atomic% or more and 86.5 atomic% or less), and the heat processing temperature is also suppressed low.

Fig. 31 shows the optimum heat treatment temperature for obtaining the maximum permeability µ 'in the soft magnetic alloy of Zr + Nb = 6 atomic%.

In particular, the permeability (μ ') is increased at the heat treatment temperature of 525 ° C. or lower because Zr is 1.5 atomic% or more and 2.5 atomic% or less, B is 8 atomic% or more and 9 atomic% or less, and the sum of Fe and Nb is 89. It turns out that it is a ribbon of atomic% or more and 90 atomic% or less (Fe is 85 atomic% or more and 86 atomic% or less).

In FIGS. 30 and 31, a ribbon of a soft magnetic alloy having Zr of 0.5 atomic% or more and 3.5 atomic% or less, B of 7 atomic% or more and 10.5 atomic% or less, and the sum of Fe and Nb of 90 atomic% or less It can be seen that when the heat treatment temperature is 550 ° C. or less, the optimum coercive force (Hc) and permeability (μ ′) are exhibited, and the composition ranges are the coercive force (Hc) and permeability (shown in FIGS. 24, 25, and 28). μ ') and the average composition range of the average grain size of the bcc-Fe phase, and if the heat treatment temperature is less than 550 ℃, the average grain size of the bcc-Fe phase is kept in a fine state and excellent soft magnetic properties It is suggested to indicate.

Next, the crystallization temperature (T _x1 ) of the bcc-Fe phase is shown in FIG. 32, the crystallization temperature (T _x2 ) of another compound phase is shown in FIG. 33, and the crystallization temperature (T _x1 ) of the other compound phase is shown in FIG. The relationship between these crystallization temperatures is T _x1 &lt; T _{x '} &lt; T _x2 ). 35 shows the interval (ΔT _x : ΔT _x = T _x2 -T _x1 ) of the crystallization temperature.

As shown in FIG. 32, it is understood that T _x1 is in the range of 464 to 500 ° C, depends on the composition ratio of Fe, Nb and Zr, and does not depend on the composition ratio of B.

Moreover, in FIG. 24 and FIG. 25, the range which shows the favorable value of coercive force (Hc) and permeability (micro ') [Zr / (Zr + Nb) is 1.5 / 6 or more and 2.5 / 6 or less, B is 8 atomic% or more 9 Atom% or less, Fe is 80 atom% or more, and the sum total of Fe and Nb is 89 atomic% or more and 90 atomic% or less], _Tx1 is the range of 480-490 degreeC.

As shown in FIG. 35, the interval ΔT _x of the crystallization temperature is in the range of 313 to 344 ° C., and the interval ΔT _x of the crystallization temperature tends to widen as the composition ratio of B decreases. In particular, when Zr is 1 atomic% or more and 2.5 atomic% or less, and B is in the range of 9.5 atomic%, it exhibits an interval (ΔT _x ) of the crystallization temperature of 330 ° C. or more. Precipitation can be suppressed and deterioration of the soft magnetic properties of the soft magnetic alloy can be prevented.

However, in FIG. 34, the plot without the subscript indicates an alloy in which the crystallization temperature (T _x1 ) of the other compound phase was not observed, and in contrast to FIGS. 24 to 26 and the like, an alloy in which T _x1 does not exist is generally better magnetic. It can be seen that it has characteristics.

As described above with reference to FIGS. 24 to 35, when the total amount of Zr and Nb is 6 atomic%, Zr is 1.5 atomic% or more and 2.5 atomic% or less, and B is 8 atomic% or more and 9 atomic% or less. , When the total amount of Fe and Nb is 89 atomic% or more and 90 atomic% or less (Fe is 85 atomic% or more and 86 atomic% or less), it shows excellent soft magnetic properties and superior lead than when Zr is 2 atomic%. It became clear that it exhibited magnetic properties.

In addition, as shown in Figs. 24 to 35, within the composition range, the precipitation of the (200) plane of Fe in the quenched state is not observed, and the ribbon becomes almost amorphous top ( Ribbons (present) exist, and a mixture of amorphous and crystalline phases ( The display) is mainly scattered outside the above range.

Thus, in the quenched state, when the ribbon mainly composed of the amorphous phase is heat-treated, it becomes clear that there is a tendency to exhibit excellent soft magnetic properties.

Specifically, high saturation magnetic flux density low iron loss Fe having a composition of Fe _85.5 Zr ₂ Nb ₄ B _8.5 , Fe ₈₅ Zr _1.75 Nb _4.25 B ₉ , Fe _85.25 Zr _1.75 Nb _4.25 B _8.75 , Fe _85.75 Zr _2.25 Nb _3.75 B _8.25 It can be seen that the soft magnetic alloy exhibits particularly excellent soft magnetic properties.

「Zr + Nb = 6.5 atomic percent soft magnetic alloy」

As shown in FIG. 36, the soft magnetic alloy whose Zr + Nb = 6.5 atomic% shows the coercive force (Hc) of 43-108 mOe.

Here, the coercive force (Hc) indicates 100 m Oe or less, the total amount of Fe and Nb is 90.5 atomic% or less, and the 50 m Oe or less indicates that Zr is 1.6 atom% or more and 3.2 atom% or less, and B It is a range of 6.75 atomic% or more and 9 atomic% or less, and the sum total of Fe and Nb is 89 atomic% or more, More preferably, it is the composition of 89.5 atomic% or more and 90.25 atomic% or less (Fe is 84.5 atomic% or more and 86.75 atomic% or less). Range of alloys.

Next, as shown in FIG. 37, the soft magnetic alloy whose Zr + Nb = 6.5 atomic% has shown the permeability (micro ') of 10500-45000 mOe.

Herein, the permeability (μ ') is 20000 or more, wherein B is 6.0 atomic% or more, more preferably 6.75 atomic% or more, and the total Fe and Nb are alloys in the composition range of 90.75 atomic% or less, and more than 30000 It is shown that B is an alloy having a composition range of 6.5 atomic% or more, preferably 6.75 atomic% or more, and the sum of Fe and Nb is 90.5 atomic% or less, preferably 90.25 atomic% or less, indicating 40000 or more. Zr is 1.25 atomic% or more and 2.5 atomic% or less, preferably 1.5 atomic% or more and 2 atomic% or less, and B is 8 atomic% or more and 9.25 atomic% or less, preferably 8.5 atomic% or more and 9 atomic% or less Moreover, the sum total of Fe and Nb is 89 atomic% or more and 90 atomic% or less, Preferably it is an alloy of the composition range of 89.5 atomic%.

Therefore, as can be seen from Figs. 36 and 37, Zr + Nb = 6.5 atomic% is satisfied, Zr is 1.5 atomic% or more and 2.5 atomic% or less, and B is 8 atomic% or more and 9 atomic% or less, preferably It is 8.5 atomic% or more and 9 atomic% or less, and the sum total of Fe and Nb is 89 atomic% or more and 90.5 atomic% or less, Preferably it is 89.5 atomic% (Fe is 85.5 atomic% or more and 85.5 atomic% or less, More preferably, Has a high magnetic permeability (μ ') and low coercive force (Hc), showing excellent soft magnetic properties.

Next, as shown in FIG. 38, Zr + Nb = 6. the soft magnetic alloy 5 at% shows a saturation magnetic flux density of 1.5 to 1.6 T (B _10).

39, the soft magnetic alloy having Zr + Nb = 6.5 atomic% exhibits residual magnetization (Br) of 0.37 to 0.97T.

The saturation magnetic flux density (B ₁₀ ) and the residual magnetization (Br) are not clearly correlated with the composition ratio, but it can be seen that all show excellent values.

Next, as shown in FIG. 40, in the soft magnetic alloy of Zr + Nb = 6.5 atomic%, it can be seen that the average grain size of the bcc-Fe phase is in the range of 9.8 to 11.5 nm, and has a fine crystalline structure.

Further, in the soft magnetic alloy, Zr + Nb = 6.5 atomic%, as shown in Figure 41, the range from -3 to 6 × 10 ^-7 × 10 ^-7 the magnetostriction constant (λ _s), 10 ^- It can be seen that ^seven small magnetostrictions are integers (λ _s ). The magnetostriction constant (λ _s ) tends to depend on the composition ratio of B, and it can be seen that the magnetostriction constant (λ _s ) becomes 0 when B is 8 atom% or more and 9 atom% or less.

Fig. 42 shows the optimum heat treatment temperature for obtaining the minimum coercive force (Hc) in the soft magnetic alloy having Zr + Nb = 6.5 atomic percent. The heat treatment temperature is in the range of 550 to 650 ° C, the composition showing low coercive force (Hc) is 550 ° C, and the heat treatment temperature is suppressed low.

43 shows the optimum heat treatment temperature for obtaining the maximum permeability µ 'in the soft magnetic alloy having Zr + Nb = 6. 5 atomic%. The heat treatment temperature here is in the range of 550-650 degreeC, the composition which shows high permeability (micro ') is 550 degreeC, and heat processing temperature is suppressed comparatively low.

Thus, it can be seen that the alloy having a composition range having good soft magnetic properties shown in FIGS. 36 and 37 exhibits excellent soft magnetic properties by heat treatment at a heat treatment temperature of 550 to 650 ° C as shown in FIGS. 42 and 43.

Next, Fig. 44 shows the crystallization temperature (T _x1 ) of the bcc-Fe phase, Fig. 45 shows the crystallization temperature (T _x1 ) of the other compound phases, and Fig. 36 shows the crystallization temperature (T _x2 ) of the other compound phases. The relationship between these crystallization temperatures is (T _x1 ) <(T _{x1 ′} ) <(T _x2 ). 37 shows the interval (ΔT _x : ΔT _x = T _x2 -T _x1 ) of the crystallization temperature.

As shown in FIG. 44, it is found that T _x1 is in the range of 488 to 511 ° C., depending on the composition of Fe, Nb, and Zr, and not on the composition ratio of B.

As shown in FIG. 47, the interval ΔT _x of the crystallization temperature is in the range of 305 to 325 ° C., and the interval ΔT _x of the crystallization temperature tends to widen as the composition ratio of B decreases. In particular, a ribbon of a soft magnetic alloy having Zr of not less than 1. 6 atomic% and not more than 2.5 atomic%, B is not less than 8. 5 atomic% and not more than 9 atomic%, and the sum of Fe and Nb is 89. 5 atomic%. , A wide crystallization temperature interval (ΔT _x ) of 320 ° C. or higher can prevent precipitation of the compound phase by precipitating only the bcc-Fe phase during heat treatment, thereby preventing deterioration of the soft magnetic properties of the soft magnetic alloy. .

However, in FIG. 45, the plot without the subscript indicates an alloy in which no crystallization temperature (T _{x1 '} ) of the other compound phase was observed, and it was found that the alloy in which T _x1' does not exist has generally good magnetic properties. Can be.

As described above with reference to FIGS. 36 to 47, when the total amount of Zr and Nb is 6.5 atomic%, Zr is 1.5 atomic% or more and 2.5 atomic% or less, and B is 6.0 atomic% or more, more preferably. Is 6.5 atomic% or more, more preferably 8.5 atomic% or more and 9 atomic% or less, and the total amount of Fe and Nb is 89 atomic% or more and 90.5 atomic% or less, more preferably 89.5 atomic% (Fe is 87.5 atoms % Or less, more preferably 87.0 atom% or less, still more preferably 84.5 atom% or more and 85 atom% or less), it is evident that excellent soft magnetic properties are exhibited.

36 to 47, within the composition range, precipitation of the (200) plane of Fe in the quenched state was not observed, and the ribbon mainly composed of the amorphous phase ( Ribbons (present) exist, and a mixture of amorphous and crystalline phases ( The display) is scattered outside the above range, and it can be seen that the amorphous upper phase exhibits better magnetic characteristics than the mixed crystalline phase.

Thus, in the quenched state, when the ribbon mainly composed of the amorphous phase is heat treated, it is evident that there is a tendency to exhibit excellent soft magnetic properties.

Specifically, high saturation magnetic flux density low iron loss Fe having a composition of Fe _84.5 Zr _1.6 Nb _4.4 B ₉ , Fe ₈₅ Zr ₂ Nb _4.5 B _8.5 , Fe _86.75 Zr ₃ Nb _3.5 B _6.75 , Fe _86.75 Zr _3.3 Nb _3.2 B _6.75 The magnetic magnetic alloy exhibits particularly excellent soft magnetic properties.

`` Zr + Nb = 7 atomic percent soft magnetic alloy ''

As shown in FIG. 48, the soft magnetic alloy whose Zr + Nb = 7 atomic% shows the coercive force (Hc) of 50-2500 mOe.

Here, the coercive force (Hc) of 200 m Oe or less means that the sum of Fe and Nb is an alloy having a composition range of 87.5 atom% or more, and the coercive force (Hc) of 100 m Oe or less indicates that B is 10 atomic%. Hereinafter, the alloy of the composition range whose total of Fe and Nb is 88. 5 atomic% or more and 92 atomic% or less.

Next, as shown in FIG. 49, the soft magnetic alloy whose Zr + Nb = 7 atomic% shows permeability (micro ') of 6000-44800.

The permeability (μ ') of 10000 or more here is an alloy of the composition range whose B is 10 atomic% or less, and the sum total of Fe and Nb is 88.5 atomic% or more.

Moreover, Zr is 4 atomic% or less, Preferably it is 3.5 atomic% or less, B is 6 atomic% or more and 9 atomic% or less, and the sum total of Fe and Nb shows 890 or more magnetic permeability (micro '). An alloy having a composition range of 5 atomic% or more, preferably 90 atomic% or more and 92 atomic% or less (Fe is 84 atomic% or more and 87 atomic% or less).

Therefore, as can be seen from FIGS. 48 and 49, Zr + Nb = 7 atomic% is satisfied, and Zr is 4 atomic% or less, preferably 3. 5 atomic% or less, and B is 6 atomic% or more and 9 atoms The ribbon of the soft magnetic alloy having a% or less, and a total of Fe and Nb of not less than 89.5 atomic%, preferably 90 or more and 92 atomic% (Fe is 84 or more and 87 atomic% or less), is 20,000 or more. It has high magnetic permeability (µ ') and low coercive force (Hc) of 100 m Oe or less, which shows excellent soft magnetic properties.

Next, as shown in FIG. 50, the soft magnetic alloy Zr + Nb = 7 atomic% shows an 1.42 to 1.68 T saturation magnetic flux density (B _10). The saturation magnetic flux density (B ₁₀ ) depends on the composition ratio of Zr and B, and when B is 9 atomic% or less, the saturation magnetic flux density (B ₁₀ ) is 1.5 T or more and B is 6 atomic% or more and 8. 5 atomic% or less, preferably represents a saturated magnetic flux density (B ₁₀₎ or more equal to or less than 8 atom% 1.55 T.

As shown in Fig. 51, the soft magnetic alloy having Zr + Nb = 7 atomic% shows residual magnetization (Br) of 0.78 to 1.44 T. Residual magnetization (Br) tends to depend on the composition ratio of B and Fe, and if B is 7 atomic% or more and 9 atomic% or less, preferably 8 atomic%, and the sum of Fe and Nb is 88.5 atomic% or less, 1.2 T The above residual magnetization (Br) is shown.

As shown in Fig. 52, in the soft magnetic alloy having Zr + Nb = 7 atomic%, it can be seen that the average grain size of the bcc-Fe phase is in the range of 9.1 to 16.7 nm.

The average grain size is almost dependent on the composition ratio of Zr and Nb, and the average grain size is 14 nm or less, the composition of Zr is 5 atomic% or less, and the average grain size is 12 nm or less, the composition of Zr is 3 atomic% or less It is understood that the average grain size of 10 nm or less is a composition of Zr of 1 atomic% or less, and these ribbons all have a fine crystalline structure.

Further, in the soft magnetic alloy, Zr + Nb = 7 atomic%, as shown in Figure 53, is in the range of -10 × 10 ^-7 to 19 × 10 ^-7 to magnetostriction constant (λ _s), 10 ^- It can be seen that a constant (λ _s ) of ^seven magnetostrictions is obtained. The magnetostriction constant (λ _s ) tends to depend on the composition ratio of B, and when B is 7.5 atomic% or more and 8.5 atomic% or less, the magnetostriction constant (λ _s ) becomes almost zero.

FIG. 54 shows the optimum heat treatment temperature for obtaining the minimum coercive force (Hc) in the soft magnetic alloy having Zr + Nb = 7 atomic%.

Among them, the coercive force (Hc) becomes smaller at the heat treatment temperature of 650 ° C. or lower, Zr is 5 atomic% or less, B is 5. 5 atomic% or more, preferably 6 atomic% or more and 11 atomic% or less, and Fe It turns out that the sum of and Nb is a ribbon of 87 atomic% or more.

Further, the coercive force (Hc) decreases at a heat treatment temperature of 600 ° C. or lower, where Zr is 2.5 atomic% or more and 3.5 atomic% or less, B is 6 atomic% or more and 8 atomic% or less, and the total of Fe and Nb is 89 atoms. It turns out that it is a ribbon of% or more and 92 atomic% or less (Fe is 85 atomic% or more and 87 atomic% or less).

In Fig. 55, the optimum heat treatment temperature for obtaining the maximum permeability µ 'in the soft magnetic alloy having Zr + Nb = 7 atomic% is shown.

Especially, it turns out that Zr is 5 atomic% or less, or the total of Fe and Nb is 92. 5 atomic% or less of ribbon when heat processing temperature becomes 650 degreeC or less.

Further, when the heat treatment temperature is 600 ° C. or less, the permeability (μ ′) is increased because Zr is at least 2.5 atomic%, preferably at least 3 atomic% at least 3. 5 atomic%, and B is at least 5. 5 atomic%, Preferably it is 6 atomic% or more and 8 atomic% or less, and the sum total of Fe and Nb is 89 atomic% or more and 91 atomic% or less (Fe is 87.5 atomic% or less, Preferably 85 atomic% or more and 87 atomic% or less) It is understood that it is a ribbon.

Next, Fig. 56 shows the crystallization temperature (T _x1 ) of the bcc-Fe phase, Fig. 57 shows the crystallization temperature (T _x2 ) of the other compound phases, and Fig. 58 shows the crystallization temperature (T _x1 ) of the other compound phases. The relationship between these crystallization temperatures is T _x1 &lt; T _{x1 '} &lt; T _x2 . 59 shows the interval (ΔT _x : ΔT _x = T _x2 -T _x1 ) of the crystallization temperature.

As shown in FIG. 56, it is found that T _x1 is in the range of 491 to 533 ° C., and depends on the composition ratio of Nb and Zr and hardly depends on the composition ratio of B.

As shown in Fig. 59, the interval ΔT _x of the crystallization temperature is in the range of 181 to 316 ° C, and the interval of the crystallization temperature is decreased as the composition ratio of Zr decreases.

(ΔT _× ) tends to be widened. The spacing of the crystallization temperature (ΔT _x) representing the more than 200 ℃ is, the Zr is 5 at% or less, and the range that the total amount of Fe and at least 87 at% Nb.

In addition, it indicates the interval (ΔT _x) of at least 300 ℃ crystallization temperature, and Zr is less than 3 atomic%, and B is 6.5 atomic% or more, preferably not more than 8 atomic% 7 atomic%, Fe and Nb When the total amount of is at least 89 atomic% (Fe is 85 atomic% or more), only the bcc-Fe phase can be precipitated during the heat treatment to suppress the precipitation of the compound phase, thereby preventing deterioration of the soft magnetic properties of the soft magnetic alloy. have.

However, in FIG. 58, the plot without a subscript indicates an alloy in which the crystallization temperature (T _{x1 '} ) of other compound phases is not observed, and it is found that an alloy in which T _x1' does not exist has generally good magnetic properties. Can be.

As described above with reference to FIGS. 48 to 59, when the total amount of Zr and Nb is 7 atomic%, Zr is 4 atomic% or less, preferably 3 atomic% or less, and B is 6 atomic% or more and 9 atoms It became clear that excellent soft magnetic properties were exhibited when it was% or less, Preferably it is 7 atomic% or more and 8 atomic% or less, and the total amount of Fe and Nb was 89 atomic% or more and 91 atomic% or less.

Specifically, high saturation magnetic flux density low iron loss Fe having a composition of Fe ₈₅ Zr ₁ Nb ₆ B ₈ , Fe ₈₅ Zr _1.2 Nb _5.8 B ₈ , Fe ₈₅ Zr ₂ Nb ₅ B ₆ , Fe ₈₆ Zr _2.4 Nb _4.6 B ₇ The magnetic magnetic alloy exhibits particularly excellent soft magnetic properties.

60 shows a soft magnetic alloy of Zr + Nb = 5. 75 atomic% and B = 8.5 atomic% (Fe _85.75 Zr _X Nb _5.75-X B _8.5 ), Zr + Nb = 6 atomic% and B = 8 to 9 atoms Soft magnetic alloys (Fe _a Zr _x Nb _6-x B _z , where a is 85 to 86, z is 8 to 9), Zr + Nb = 6. 25 atomic% and B = 8. 25 atomic% One soft magnetic alloy (Fe _85.5 Zr _x Nb _6.25-x B _8.25 ), Zr + Nb = 6. 5 atomic% and B = 8. 5 atomic% soft magnetic alloy (Fe ₈₅ Zr _x Nb _6.25-x B _8.5 ), Zr + Nb = 7 atomic% and B = 8 to 9 atomic% in the soft magnetic alloy (Fe _a Zr _x Nb _7-x B _z , where a is 84 to 85, z is 8 to 9), The relationship between the composition ratio of Zr and Nb and the coercive force (Hc) is shown.

61 shows the relationship between the composition ratio of Zr and Nb and the interval (ΔT _x : ΔT _x = T _x2 -T _x1 or ΔT _x = T _{x1 '} -T _x1 ) in the alloy.

In FIG. 60, in the range where Zr / (Zr + Nb) is 0 to 0.4, all alloys have a coercive force (Hc) of 0.1 Oe or less, but when Zr / (Zr + Nb) is larger than 0.5, the coercive force (Hc) increases. It can be seen that the soft magnetic properties deteriorate.

In addition, Zr + Nb = 6 atomic% of the alloy has a Zr / (Zr + Nb) of more than 0.1 and less than 0.5 to obtain a coercive force (Hc) less than Zr / (Zr + Nb) = 0, composite Zr and Nb It can be seen that the soft magnetic properties improve when added. However, if Zr / (Zr + Nb) exceeds 0.5, the coercive force Hc is rather deteriorated, which is not preferable.

In Fig. 61, Zr / (Zr + Nb) is in the range of 0 to 0.7, particularly in the range of more than 0 and 0.4 or less, all alloys are in the range of 200 ° C or more (Zr / (Zr + Nb) is more than 0 and 0.4 or less. At 300 ° C. or higher), the interval (ΔT _x ) of the crystallization temperature is higher than Zr / (Zr + Nb), but exceeds 0.7, and the interval (ΔT _x ) of the crystallization temperature is rapidly narrowed.

As described above, when the value of Zr / (Zr + Nb) increases, the interval ΔT _x of the crystallization temperature is narrowed, and compound phases other than the bcc-Fe phase are easily precipitated at the time of heat treatment, and the coercive force (Hc) is considered to increase.

Example 3 Regarding Fracture Strain (λ _f )

In the same manner as in Example 1, a ribbon having a sheet thickness of 20 μm was prepared using Fe _85.5 Zr ₂ Nb ₄ B _8.5 , Fe ₉₀ Zr ₇ B ₃ , and Fe ₈₄ Nb ₇ B ₉ . Next, a heat treatment was performed by heating from 510 ° C to 670 ° C at a heating rate of 180 ° C / min for 5 minutes to prepare a ribbon of soft magnetic alloy.

The fracture strain (λ _f ) was measured for the ribbon after the obtained heat treatment. The fracture strain (λ _r ) was obtained from the ribbon bending radius when the ribbon was bent and separated. The results are shown in FIG. 62.

As can be seen from FIG. 62, the fracture strain (λ _f ) of the ribbon of Fe _85.5 Zr ₂ Nb ₄ B _{8.5 according} to the present invention, in which the composition ratio of Zr and Nb was adjusted and added at the same time, was subjected to heat treatment at 510 ° C. When 12.71 x 10 ^-3 was shown and heat-treated at 520 ° C, 11.98 x 10 ^-3 was shown, indicating that the fracture strain (λ _r ) was superior to 10 x 10 ^-3 and had excellent workability.

On the other hand, a ribbon having a composition of Fe ₉₀ Zr ₇ B ₃ exhibits a breakage strain (λ _f ) of 8.35 × 10 ⁻³ at the time of heat treatment at 620 ° C., and a ribbon having a composition of Fe ₈₄ Nb ₇ B _{9 has} a composition at 630 ° C. The breakdown strain (λ _f ) of 9.72 × 10 ⁻³ is exhibited at the time of heat treatment, but not all exceed 10 × 10 ⁻³ .

Example 4 Regarding Iron Loss

_{Rapid cooling} of the composition of Fe ₈₅ Zr _1.75 Nb _4.25 B ₉ , Fe _85.5 Zr ₂ Nb ₄ B _8.5 , Fe _85.75 Zr _2.25 Nb _3.75 B _8.25 , Fe ₇₈ Si ₉ B ₁₃ (commercially available: amorphous alloy) in the same manner as in Example 1. The ribbon of the state was produced. However, for Fe ₈₅ Zr _1.75 Nb _4.25 B ₉ , the melt temperature is 1260 to 1280 ° C., and for Fe _85.75 Zr _2.25 Nb _3.75 B _8.25 and Fe _85.5 Zr ₂ Nb ₄ B _8.5 , the temperature of the melt is 1300 ° C. It was set as. The ribbon thus obtained was wound to a ring shape having an outer diameter of 10 mm and an inner diameter of 6 mm, and laminated thereon to form a core. Next, the iron loss was measured by heating at 510 ° C. to 525 ° C. at a temperature increase rate of 180 ° C./minute to maintain the heat for 5 minutes.

Fig. 63 shows the result of iron loss measured while applying a magnetic flux having a frequency of 50 Hz at a room temperature to a core made of a ribbon obtained by heat treatment at 510 ° C or 520 ° C.

As can be seen in Figure 63 the core of the composition of Fe ₈₅ Zr _1.75 Nb _4.25 B ₉ , Fe _85.5 Zr ₂ Nb ₄ B _8.5 , Fe _85.75 Zr _2.25 Nb _3.75 B _{8.25 according} to the present invention is Fe ₇₈ Si ₉ B ₁₃ It is evident that iron loss is lower than that of the core of the composition, and in particular, even when the magnetic flux density (Bm) is 1.4 T, the iron loss is 0.1 W / kg or less.

64 shows iron loss before and after heat treatment (aging) at a heating temperature of 200 ° C. and a heating time (t) of 500 hours for a core having a composition of Fe ₈₅ Zr _1.75 Nb _4.25 B ₉ and Fe ₇₈ Si ₉ B ₁₃ . And the magnetic flux density (Bm).

Cores of Fe ₇₈ Si ₉ B ₁₃ composition showed little change in iron loss before and after heat treatment, whereas cores of Fe ₈₅ Zr _1.75 Nb _4.25 B ₉ were smaller after heat treatment in areas where Bm exceeded 1.4 T. You can see that you are losing. Therefore, it can be seen that the alloy of the present invention has excellent thermal stability in the region of high Bm value.

In addition, in Fig. 65, the heating temperature in the nitrogen atmosphere is 200 for a core having a composition of Fe ₈₅ Zr _1.75 Nb _4.25 B ₉ , Fe _85.5 Zr ₂ Nb ₄ B _8.5 , Fe _85.75 Zr _2.25 Nb _3.75 B _8.25 , Fe ₇₈ Si ₉ B ₁₃ . The relationship between iron loss and heating time (t) when heated at 0 ° C. and heating time (t) from 0 to 500 hours is shown. The iron loss was measured by applying a magnetic flux density (Bm) of 1.4 T with a frequency of 50 Hz at room temperature. In addition, the coercive force (Hc) and permeability (μ ') of the core were measured. The results are shown in Tables 2 and 3. In addition, the value of the iron loss (Pcm) shown in FIG. 65 is also shown to Tables 2 and 3 simultaneously.

As can be seen from Figure 65, the core of the composition in the embodiment of the present invention can be seen that the iron loss does not change almost regardless of the increase in the heating time, it is excellent in thermal stability.

As shown in Tables 2 and 3, before and after heating, the core of the present invention tends to have a slight decrease in the permeability (μ ') and the coercive force (Hc) tends to increase. You can see that it is very small. On the contrary, it can be seen that the iron alloy of Fe ₇₈ Si ₉ B _{13 as} a comparative example has a large iron loss, and the iron loss is greatly changed with the increase of the heating time.

FIG. 66 shows the relationship between the iron loss change rate and the magnetic flux density (Bm) of each core when heated for 500 hours based on the iron loss of the core before heating in FIG. 65. Tables 4 and 5 show cores before and after heat treatment. The result of measuring the saturation magnetic flux density (B ₁₀ ), the residual magnetization (Br), the coercive force (Hc) and the magnetic permeability (μ ') is shown, and the actual measured value of iron loss (Pcm) is also shown at the same time.

A core having a composition of Fe ₇₈ Si ₉ B ₁₃ (comparative example) tends to have a small decrease in iron loss rate (change in the direction of decreasing iron loss) with an increase in magnetic flux density (Bm). On the other hand, it can be seen that the core loss rate of Fe ₈₅ Zr _1.75 Nb _4.25 B ₉ and Fe _85.5 Zr ₂ Nb ₄ B _{8.5 according} to the present invention is significantly reduced at particularly high magnetic flux density (Bm).

Further, Tables 4 and the heating before and after as shown in the fifth embodiment the core of the present invention, the permeability (μ ') tends to be a coercive force (Hc) increases slightly decreases, but the decrease is very small, Fe ₇₈ Si ₉ The core having a composition of B ₁₃ has a small permeability (μ ') and a decrease of 26% before and after heating, so that the thermal stability of the permeability (μ') is higher than that of the highly saturated magnetic flux density low iron loss Fe-based soft magnetic alloy according to the present invention. It is clear that this is falling.

In addition, in FIG. 67, the core loss of the cores of Fe _85.5 Zr ₂ Nb ₄ B _8.5 and Fe ₇₈ Si ₉ B ₁₃ was extremely high at a heating temperature of 320 ° C. and heat-treated at a heating time (t) of 0 to 100 hours. And the heating time t are shown. In addition, iron loss was measured by applying a magnetic flux of 1.4 T with a frequency of 50 Hz at room temperature.

As can be seen in Figure 60, it can be seen that the core made of an amorphous alloy of Fe ₇₈ Si ₉ B ₁₃ composition is increasing the iron loss with increasing the heating time.

On the other hand, it can be seen that the core loss of Fe _85.5 Zr ₂ Nb ₄ B _8.5 hardly changes even when the heating time increases.

68 shows the relationship between the iron loss change rate and the heating time when the iron loss of the core before heating in FIG. 67 is used as a reference. It is evident that the iron loss change rate of the core of Fe ₇₈ Si ₉ B ₁₃ phosphorus increases significantly with increasing heating time compared to the core of Fe _85.5 Zr ₂ Nb ₄ B _8.5 .

As described above, the composition of the core of Fe ₈₅ Zr _1.75 Nb _4.25 B ₉ , Fe _85.5 Zr ₂ Nb ₄ B _8.5 , Fe _85.75 Zr _2.25 Nb _3.75 B _{8.25 according} to the present invention can be _reduced even after heat treatment at 200 to 320 ° C. Since the increase of is suppressed low, the iron loss change rate is small and thermal stability is excellent.

Example 5 Addition of Zn

Next, the Example which added Zn to the composition of FeMB system or FeZrNbB system is demonstrated below.

The soft magnetic alloy ribbon sample shown in each of the following examples was created by the single-roll liquid quenching method using the apparatus shown in FIG. That is, the molten metal of the predetermined | prescribed component was blown on the said cooling roll through the quartz nozzle by the pressure of argon gas, and it quenched by the nozzle placed on one rotating copper cooling roll, and the ribbon was obtained. The width of the ribbon created as above was about 15 mm and the thickness was about 20 micrometers. The ribbon in the quenched state is composed of an alloy mainly composed of amorphous materials, but in order to precipitate fine grains of bccFe and improve soft magnetic properties, the ribbon is heated to a temperature above the crystallization temperature and then cooled to perform annealing treatment. A soft magnetic alloy ribbon sample of magnetic flux density low iron loss Fe group and each comparative sample were obtained.

The magnetic permeability of the soft magnetic alloy ribbon sample obtained as described above was measured by using an impedance analyzer after processing the ribbon into a ring shape having an outer diameter of 10 mm and an inner diameter of 6 mm, winding it on a laminate. The measurement conditions of the magnetic permeability (μ ') were 5 m Oe and 1 kPa. Coercive force (Hc) and magnetic flux density (B ₁₀ ) were measured at 10 Oe by a direct current BH loop tracer. In addition, B ₁₀ is a value almost equal to the saturation magnetic flux density (Bs).

FIG. 69 shows the structural changes of the soft magnetic alloy ribbons (Fe _0.8575 Zr _0.02 Nb _0.04 B _0.0825 ) _99.88 Zn _0.12 before and after heat treatment by X-ray diffraction analysis. In FIG. 69, in the quenched state (a state in which the alloy melt is quenched and turned into a ribbon), an amorphous, distinctive and broad diffraction figure is recognized, and after heat treatment, a diffraction figure unique to Fe of body centered cubic (bcc) is recognized, It can be seen that the structure of the alloy changed from amorphous to body-centered cubic crystal by heat treatment.

70 shows the coercive force of a sample having a composition of Fe _b Zr _d Nb _e B _x of a composition similar to that of the present invention, and a composition system in which Zn is added in the range of 0.034 to 0.142 atomic% with respect to the alloy sample of the composition; of a _{(Fe c / 100 Zr d /} 100 Nb e / 100 B f / 100) 100-z Zn triangular joseongdo showing a sample result of measuring the coercive force in the _z composition, respectively.

In the addition of Zn sample, the sample showing the numeral ① in Fig. 70 is _{(Fe 0.855 Zr 0.02 Nb 0.04 B} 0.085) 99.944 Zn sample, the sample shown in ② of the composition _0.056 is _{(Fe 0.855 Zr 0.02 Nb 0.04 B} 0.085) A sample having a composition of _98.892 Zn _0.108 , (3) is (Fe _0.855 Zr _0.02 Nb _0.04 B _0.085 ) _99.859 Zn a sample of composition _0.141 , and a sample (4) is (Fe _0.8575 Zr _0.02 Nb _0.04 B _0.0825 ) _99.96 Zn _0.04 The sample of composition, the sample shown in (⑤) is a sample of the composition (Fe _0.8575 Zr _0.02 Nb _0.04 B _0.0825 ) _99.875 Zn _0.125 , the sample shown in (⑥) is a sample of the composition of (Fe _0.8575 Zr _0.02 Nb _0.04 B _0.0825 ) _99.875 Zn _0.133 , The sample shown in (7) is (Fe _0.86 Zr _0.02 Nb _0.04 B _0.08 ) the sample of composition _99.866 Zn _0.034 , the sample shown in (8) is (Fe _0.86 Zr _0.02 Nb _0.04 B _0.08 ) the sample of composition _99.883 Zn _0.117 , the sample shown in (9) Is a sample having a composition of (Fe _0.86 Zr _0.02 Nb _0.04 B _0.08 ) _99.858 Zn _0.142 .

In Fig. 70, in the FeZrNbB-based alloy sample containing 6 atomic% of Zr and Nb in total, it can be seen that B exhibits a low coercive force below 50 m Oe by setting it to 6 to 9.5 atomic% within the range of 5 to 12.5 atomic%. Among these ranges, the coercive force was particularly low below 40 m Oe in the range of 8 to 9.5 atomic%, Zn 1.5 to 2.5 atomic%, and Fe + Nb 89 to 90 atomic%. In addition, in the samples of 1 to 9 where Zn was further added to the samples having these compositions, the low coercive force below 100 m Oe was shown. Moreover, it turned out that the value of coercive force tends to fall especially when Zn is added with respect to the sample of about 40 mOe and about 50 mOe.

In addition, in FIG. The sample indicated by Δ mark is a sample in which diffraction occurs at (200) peak of bccFe due to partial precipitation of bccFe grains when X-ray observation of the ribbon sample in the state of obtaining the ribbon sample by rapid cooling. In addition , The sample shown by the mark is a sample in which the diffraction peak in the crystal phase does not develop when X-ray observation is performed in a state where a ribbon is obtained by quenching, which means that it is completely amorphous. The magnetic properties of each of these samples show that the coercive force of the samples completely amorphous after quenching is lowered.

Fig. 71 is a triangular composition diagram showing the result of measuring the permeability (μ ': real part of complex permeability) at 1 kHz with respect to the sample described above. In the samples having the present composition and added Zn in the range of 0.034 to 0.142 atomic%, all exhibited excellent permeability of more than 30000, and the samples in which Zn was added in the range of 0.04 to 0.142 atomic% showed more than 40000.

Fig. 72 is a triangular composition diagram showing the saturation magnetic flux density (B ₁₀ ) obtained from the magnetization curve obtained by applying the applied magnetic field 10 Oe, and Fig. 73 is a triangular composition diagram showing the measurement result of the residual magnetic flux density (Br) of the previous sample. .

If the amount of Zr and Nb in the composition system similar to the composition system of the present invention, it is clear that a high saturation magnetic flux density of more than 1.5 T can be obtained. Among the composition systems, Zn is in the range of 0.034 to 0.142 atomic% in the alloy sample of the composition system over 1.6 T. All of the samples 1 to 9 added showed excellent saturation magnetic flux density over 1.6 T. Therefore, even if Zn is added within the scope of the present invention, it is clear that the saturation magnetic flux density hardly changes and maintains a high value.

FIG. 74 is a triangular composition diagram showing the first crystallization temperature (T _x1 is the crystallization temperature of bccFe) of the previous sample, and FIG. 75 is a triangular composition showing the intermediate crystallization temperature (T _x1 'is the crystallization temperature of the compound phase) of the previous sample. FIG. 76 is a triangular composition diagram showing the second crystallization temperature (T _x2 is the crystallization temperature of the compound phase) of the previous sample, and FIG. 77 is a triangular composition diagram showing ΔT _x represented by T _x2 -T _x1 .

Below, these 1st crystallization temperature, intermediate crystallization temperature, and 2nd crystallization temperature are demonstrated.

In the alloy of the composition of the present invention, when an alloy mainly composed of an amorphous phase generated by quenching is heated up, first, an exothermic reaction accompanying crystallization of the bccFe phase occurs, and crystallization of other compound phases (Fe ₃ B or Fe ₂ B, etc.) at regular intervals occurs. Exothermic reaction occurs, and another exothermic reaction occurs depending on the composition between them. The first exothermic peak is the largest exothermic peak that accompanies the crystallization of bccFe and corresponds to the first crystallization temperature, the second exothermic peak is the small exothermic peak that generates the compound phase and corresponds to the intermediate crystallization temperature, and the third exothermic peak. Is a small exothermic peak that produces another compound phase and corresponds to the second crystallization temperature. However, the 2nd exothermic peak may not express depending on a composition, and the sample of-mark shown in FIG. 75 is a sample in which the 2nd exothermic peak did not express (the sample in which the intermediate crystallization temperature _Tx1 'did not express). Moreover, the composition which does not express a 2nd exothermic peak is excellent in a magnetic characteristic. It can be seen that there is almost no change in crystallization temperature even if Zn is added to these composition systems.

In this way it is preferable that the interval (ΔT _x) of the crystallization temperature determined by more than 200 ℃. Although (ΔT _x ) shown in FIG. 77 is all 200 ° C. or more, when the temperature is 200 ° C. or more, the crystallization temperature interval between the bccFe phase and the compound phase becomes wider, so that it is easier to heat-treat the alloy under optimum conditions, thereby depositing only the bccFe phase and other compounds. By suppressing the precipitation of the soft magnetic properties can be easily improved. Therefore, the heat treatment temperature of the alloy is preferably carried out between the first crystallization temperature and the second crystallization temperature (temperature between T _x1 and T _x2 ).

FIG. 78 is a triangular composition diagram showing the grain size of a composition sample which does not contain Zn in a composition system similar to the composition of the present invention. However, when the Zn of the composition range of the present invention is added to this composition system, the grain size decreases slightly. The inventors confirm. Therefore, it can be seen that even in the alloy of the composition system of the present invention, a particle size of 12 nm or less, preferably 11 nm or less, can be obtained.

Fig. 79 is a triangular composition diagram showing the magnetostriction (? S) of a composition-based sample containing no Zn in a composition system similar to the composition of the present invention, but the present inventors have confirmed that the magnetostriction is equivalent even if Zn is added to the composition system. Accordingly, it can be seen that the magnetostriction is near 0 even in the alloy of the composition system of the present invention in which Zn is added to the alloy of the composition system shown in FIG. 79.

Fig. 80 shows the Zn concentration dependency in grain size (D) of the inventive composition-based alloy sample containing Zn. The grain size was slightly but decreased by the effect of Zn addition.

Fig. 81 shows the Zn concentration dependence of (? S) on the magnetostriction of the inventive composition-based alloy sample containing Zn. The effect of Zn addition clearly shows a tendency to decrease the magnetostriction, but the amount of change is slight.

82 is a Fe _85.75 Zr ₂ Nb ₄ B _8.25 of the measurement results of the iron loss in AC magnetization characteristic measuring apparatus with respect to Zn in the alloy samples of the composition to the sample was added 0.12 at% or 0.13 atomic percent of the comparative example Fe ₇₈ Si ₉ B _{It is} shown in comparison with the numerical value of the ribbon sample of ₁₃ compositions. As can be seen from the results shown in FIG. 81, it was evident that the iron loss of the sample of the present invention showed a smaller iron loss than the sample of the comparative example. It can be seen that the sample of the present invention is less than 0.1 W / kg at 1.5 T in iron loss, and is an excellent value of about 1/10 of a silicon steel sheet and an excellent value of a few minutes of Fe group amorphous.

83 shows a comparative example sample of Fe ₇₈ Si ₉ B ₁₃ phosphorus composition, a comparative example sample of Fe ₈₅ Zr _1.75 Nb _4.25 B ₉ phosphorus composition, a comparative example sample of Fe _85.5 Zr ₂ Nb ₄ B _8.5 phosphorus composition, Fe _85.75 Zr _2.25 Comparative Example Sample of Nb _3.75 B _8.25 Phosphorus Composition and (Fe _0.8575 Zr _0.02 Nb _0.04 B _0.0825 ) _99.88 Zn The temporal change of the iron loss of the alloy sample of the composition of _0.12 phosphorus (heated to 200 ° C. for a predetermined time and then returned to room temperature Measurement).

As can be seen from the results shown in FIG. 83, the sample of the present invention exhibited excellent characteristics with almost no time change in the state of iron loss much smaller than that of the comparative example of the composition of Fe ₇₈ Si ₉ B ₁₃ . Moreover, it turns out that it is lower iron loss and a change rate is low compared with the comparative example sample of the substantially equivalent composition which does not add Zn. Moreover, the sample of this example shows the outstanding value which iron loss is less than 0.1 W / kg after 300 hours of heating.

5 and 6 show the elapsed time dependence of iron loss, coercive force, and permeability of each sample used in FIG. 83. As can be seen from the results shown in Tables 1 and 2, the iron loss value itself (0.081 to 0.90) of the sample according to the present invention to which Zn was added was low, the rate of change was low, the coercive force was small (0.038), and the magnetic permeability was also low. It can be seen that it is excellent (60200 to 61200). In addition, the samples containing Zr and Nb in the prescribed range have low iron loss and coercive force and high permeability, but it can be seen that addition of Zn to the composition system results in even lower iron loss and higher permeability.

FIG. 84 shows the result of measuring iron loss at room temperature after heating at 320 ° C. for a predetermined time using a sample having the same composition as that used in FIG. 83, and FIG. 85 shows the rate of change of iron loss shown in FIG. 84.

From the results shown in FIGS. 84 and 85, the sample of the present invention showed a much lower iron loss rate than the comparative sample of the composition of Fe ₇₈ Si ₉ B ₁₃ , and the composition of Fe _85.75 Zr _2.25 Nb _3.75 B _8.25 phosphorus containing no Zn It became clear that the iron loss rate was lower than that of the comparative sample.

From these, it was found that by adding a small amount of Zn to the FeNbZrB-based alloy in the range stipulated in the present invention, the iron loss can be made lower, and the rate of change of iron loss over time can also be lowered.

86 is a ribbon sample having a sheet thickness of 20 μm, respectively, and a comparative example sample having a composition of Fe ₇₈ Si ₉ B ₁₃ , a comparative example sample having a composition of Fe ₈₄ Zr _3.5 Nb _3.5 B ₈ Cu ₁ , and Fe ₉₀ Zr ₇ B A comparative example sample of the composition of ₃ phosphorus, a comparative example sample of the composition of Fe ₈₄ Nb ₇ B ₉ , a comparative example sample of the composition of Fe _73.5 Si _13.5 B ₉ Nb ₃ Cu ₁ , and Fe _85.5 Zr ₂ Nb ₄ B ₈ phosphorus Comparative Example of Composition With respect to the sample and the alloy sample of the present invention having a composition of (Fe _0.855 Zr _0.02 Nb _0.04 B _0.085 ) _99.86 Zn _0.14 , whether the bending diameter (Df: mm: bending degree without breaking up to what bending radius was possible) ) And fracture strain (λf: 10 ^-3 : strain upon fracture).

In this case, the bending deformation is to use two rods and a ribbon sample, sandwich the rods arranged between the two rods, and gradually approach the two rods to bend the ribbon into a mountain shape. When the ribbon was bent, the distance between end surfaces of the rod when the ribbon was folded and cut was L, and when the thickness of the ribbon was t, the value of t / (Lt) was defined as fracture strain (λf). The result is shown in FIG.

In the sample of the present invention having the composition shown in Fig. 86 to Fig. 86, the bending diameter can be made small at an appropriate heat treatment temperature of 510 to 520 캜, and it was difficult to be broken. Since the heat treatment temperature is different from the crystallization temperature in each composition system, the temperature is different as shown in FIG. 86. The temperature increase rate during heat treatment is 180 deg. C / min for each sample, and is maintained for 5 minutes at the prescribed heat treatment temperature. It was set as the process of cooling later.

The excellent fracture bending characteristics are effective in the case that the ribbon of soft magnetic alloy is wound to form the magnetic core of the transformer, and the ribbon is not cracked. The smaller the bending diameter, the smaller the bending radius of the ribbon is. Means that.

Fig. 87 shows the Zn concentration dependence at the amorphous Curie temperature determined by the temperature change of magnetization, and Fig. 88 shows the Zn concentration dependence at the Curie temperature of the alloy in the quenched state without heat treatment.

In the sample which was not heat-treated, the change of Curie temperature accompanying the change of Zn concentration is not observed, but it is thought that this is because Zn concentration is low. On the other hand, after heat treatment at 510 ° C., the Curie temperature is increased with the increase of Zn concentration. This is assumed to be due to the precipitation of the bccFe phase and the change in the amorphous phase structure of the remainder.

The present inventors found that Zn was concentrated in the residual amorphous phase as a result of TEM observation and compositional analysis of these samples. It is thought that this thickening of Zn raises the Curie temperature of the amorphous phase. The inventors assume that raising the Curie temperature of the residual amorphous phase is linked to raising the exchange bonds between the bccFe phases, leading to an increase in permeability and a decrease in the coercive force.

FIG. 89 shows the Zn concentration dependence in the coercive force of a soft magnetic alloy having a composition of (Fe _0.86 Nb _0.07 B _0.07 ) _100-z Zn _z , and it is clear that the coercivity tends to decrease by the addition of Zn, and the Zn concentration The coercivity became the minimum level in the range of 0.04 to 0.07 atomic%, and the coercivity gradually increased with the increase of Zn concentration, but the coercivity of 0.12 atomic% showed a lower coercive force than the sample to which Zn was not added.

Fig. 90 shows the Zn concentration dependence in the magnetic permeability of the soft magnetic alloy of the same composition. By adding Zn, the permeability was improved, the maximum at 0.07 atomic%, and then the permeability gradually decreased.

FIG. 91 shows the Zn concentration dependence in the coercive force of a soft magnetic alloy having a composition of (Fe _0.86 Zr _0.02 Nb _0.04 B _0.08 ) _100-z Zn _z , similar to the test results of the alloy shown in FIG. 89. Since the coercive force has a minimum value and the sample added with 0.133 atomic% Zn exhibits a lower value of about 65% than the sample without the Zn added, it became clear that the coercive force can be reduced by adding Zn.

Fig. 92 shows the Zn concentration dependence in the magnetic permeability of the soft magnetic alloy of the same composition. It is clear that the permeability is improved by adding Zn, and the permeability is shown to be the maximum by adding 0.133 atomic%. Moreover, about the Zn addition amount, the sample added with 0.025 atomic% showed the high value of 29821, and also showed the excellent permeability exceeding 31769 also in the sample which added 0.19 atomic%. From these, it is understood that in order to obtain a high permeability of 30000 or more, it is important to add Zn in an amount exceeding 0.025 atomic%, and it is important to add Zn in an amount of 0.2 atomic% or less.

The test results of the coercive force and the magnetic permeability shown in FIGS. 91 and 92 are superior to those of the test results shown in FIGS. 89 and 90 in which the content ratio of Zr and Nb of the sample used in each test tends to be good in soft magnetic properties. Because it was.

Table 7 shows the permeability (μ ': 1 kHz), coercive force (Hc: Oe) and the coercive force when Zn is added to the FeNbB alloy sample, FeZrB alloy sample, FeHfB alloy sample and FeZrNbB alloy sample. The value of saturation magnetic flux density (B ₁₀ : T) was described as an example. Moreover, about the sample which does not add Zn, it shows together in Table 3 as a comparative example.

In Table 7, Z7 was added in a sample of FeNbB-based No. 16 to 0.07 atomic%, and Zn was added in a sample of FeZrB-based No.17 to 0.1 atomic%. Zn was added to the sample of system No. 18 by 0.1 atomic%, and Zn was added to the sample of No. 19 by 0.13 atomic%. The sample of No. 14 was added with 0.13 atomic% of Zn, and the sample of No. 15 was added with 0.14 atomic% of Zn to the sample of No. 21.

From the results shown in Table 7, it was clear that the addition of Zn in any composition system markedly improved permeability, decreased coercive force, and saturation magnetic flux density exhibited excellent values around 1.6T.

By the way, when Zr and Nb are added as an element (M), the range of the addition amount and the value of Zr / (Zr + Nb) is the same as the FeZrNoB type alloy mentioned above. However, as shown in FIG. 93, since the coercive force becomes very high when the addition amount of Zr + Nb is 4 atomic%, the lower limit of the amount of Zr + Nb is estimated to be 5 atomic% in any of the FeZrNbB and FeZrNbBZn-based alloy systems.

When the soft magnetic alloys described above are used as magnetic cores, the loss can be reduced, the size can be reduced, and the low iron loss magnetic cores excellent in workability can be obtained.

Example 6 First Manufacturing Method

The highly saturated magnetic flux density low iron loss Fe-based soft magnetic alloy shown in each of the following examples was prepared by the single-roll liquid quenching method using the apparatus shown in FIG. That is, molten metal or melt of a predetermined component was ejected by the pressure of Ar gas from the quartz nozzle placed on one rotating steel cooling roll, and quenched to obtain a ribbon. At this time, the temperature of the melt ejected from the nozzle, that is, the injection temperature, is about 1240 to 1350 ° C in this embodiment. It is needless to say that the above-mentioned "molten metal of a predetermined component" corresponds to the highly saturated magnetic flux density low iron loss Fe-based soft magnetic alloys having the respective compositions in the first to third cases described above. However, the following caution is required for the second and third cases containing Zn. That is, considering that the melting point and boiling point of Zn are 419.5 ° C and 908 ° C, respectively, and the injection temperature is about 1240 to 1350 ° C, most of Zn is evaporated and lost. Therefore, in the case where the composition ratio of Zn is to be finally secured as an alloy that satisfies the composition ratio in the second and third cases, it is necessary to inject Zn in an amount exceeding the above Zn amount into the crucible. Specifically, in order to make the target composition, that is, 0.025 atomic% or more and 0.2 atomic% or less, the raw material Zn of about 20 times may be added to the crucible.

The width of the ribbon created as above was about 15 mm and the thickness was about 20 micrometers. The quenched ribbon is composed of an alloy mainly composed of amorphous materials. As described above, in order to precipitate fine crystal grains of bccFe and improve soft magnetic properties, the predetermined heat treatment temperature Ta or higher than the crystallization temperature (T _x1 ) Heat treatment was performed after the furnace was cooled and cooled to obtain a highly saturated magnetic flux density low iron loss Fe-based soft magnetic alloy ribbon sample and each comparative sample. At this time, in the present embodiment, the heat treatment temperature Ta is defined to be 490 ° C. or higher and 670 ° C., and the speed of reaching the heat treatment temperature Ta, that is, the temperature increase rate is 10 ° C./min or more and 200 ° C./min or less. It was prescribed. However, in order to obtain samples, such as a comparative example, the alloy by the conditions which changed the holding time after reaching | attaining heat processing temperature (Ta), etc. was also created.

The magnetic permeability (μ ') of the soft magnetic alloy ribbon sample obtained as described above (the real part of the complex permeability) is obtained by processing a ribbon to form a ring shape having an outer diameter of 10 mm and an inner diameter of 6 mm, winding this on an overlapping impedance analyzer, and Measured using. The measurement conditions of the magnetic permeability (μ ') were 5 m Oe and 1 KHz. The coercive force (Hc) and the saturation magnetic flux density (B ₁₀ ) were measured by measurement 10 Oe using a direct current BH loop tracer.

Hereinafter, a description will be made focusing on how the change in the heat treatment temperature (Ta) and the temperature increase rate of the sample obtained as described above affects all the properties of the highly saturated magnetic flux density low iron loss Fe-based soft magnetic alloy.

94 and 95 show that the heat treatment temperature Ta and the temperature increase rate of the alloy having the composition Fe _85.5 Zr ₂ Nb ₄ B _8.5 was subjected to heat treatment in which the holding time after reaching the heat treatment temperature Ta reached 0 min. It is shown about the result of having measured what is the relationship between permeability (micro ') (real part of complex permeability) (FIG. 94) and coercive force (Hc) (FIG. 95) in an alloy. In these figures, the thick solid line indicates that the injection temperature is 1280 占 폚 and the thin solid line indicates that the injection temperature is 1320 占 폚, respectively. 4 to 11 which will be described later are the same as the above-mentioned drawings, and these are different from those shown in FIGS. 94 and 95 only in the holding time at the heat treatment temperature Ta. Moreover, although it is about the composition of the said alloy, it turns out that this satisfy | fills the conditions regarding the said 1st alloy.

First, the permeability (μ ') of FIG. 94 is shown, and as can be seen from the figure, it turns out that it distributes to 200-53800 as a whole. Among these, the parts which can be said to be particularly good are those where the heat treatment temperature (Ta) is about 520 to 550 ° C and the temperature increase rate is about 40 ° C / min, and the heat treatment temperature (Ta) is about 650 ° C and the temperature increase rate is 40 It can be said that it crosses about degree C / min. In this part, the permeability (μ ') is 50000 to 53800 without affecting the injection temperature, and it can be seen that a very good value is measured. Moreover, permeability (micro ') = 40000 is recorded also in the vicinity, and this is also favorable.

In addition, the dotted line indicated by "amor." In the figure indicates that all of the samples were amorphous in the region where the heat treatment temperature Ta was low among the two regions divided by the dotted line, that is, the microcrystalline grains of bccFe described above did not crystallize. It suggests what it did not. In this case, as is clear from the figure, it can be seen that the permeability µ 'is measured only at a very low value. As a result, it can be said that the importance of the bccFe phase in the present invention is recognized.

On the other hand, about the coercive force Hc of FIG. 95, it turns out that it is observed in the range of 39-2455 m Oe. Particularly favorable as the soft magnetic properties is the place where the heat treatment temperature Ta crosses around 550 ° C and the temperature increase rate is 100 ° C / min. At this point the coercive force (Hc) is 39 m Oe. From this figure, it can be said that 50 m Oe is observed in most of the regions, which is generally good. As in the case of the above magnetic permeability (μ '), the areas where the entire sample is in an amorphous phase are white or thin. Order coercive force (Hc) of the cloth is measured, and it is not said to be preferable.

Next, FIG. 96 and FIG. 97 and subsequent description. 96 and 97 show results of whether the magnetic permeability μ 'and the coercive force Hc were measured as in ° C when the holding time after reaching the heat treatment temperature Ta was sequentially changed as described above. Specifically, FIGS. 96 and 97 have a holding time of 5 min, hereinafter, FIGS. 98 and 99 are 10 min, 100 and 101 are 30 min, and 102 and 103 are 60 min, respectively. .

In the magnetic permeability (μ ') of FIG. 96, when injection temperature is 1320 degreeC, heat processing temperature (Ta) is 500-600 degreeC, and a temperature rising rate crosses about 30-90 degreeC / min, and injection temperature is 1280 degreeC Has a permeability (μ ') of about 50000 in an area slightly narrower than that defined by the above.

First, in FIG. 97, the coercive force Hc shows about 40 mOe at the injection temperature of 1320 ° C where the heat treatment temperature Ta intersects around 530 ° C and the temperature increase rate is about 40 to 100 ° C / min. When the injection temperature is 1280 ° C, the optimum region is slightly narrowed similarly to the magnetic permeability μ 'in FIG. 96.

By the way, it is noted in FIG. 96 and FIG. 97 that the above-mentioned area of "amor." Has disappeared. This is the same also after FIG. That is, it can be surmised that the above-mentioned "amor." Region was formed because the energy sufficient for the growth of the bccFe crystals was not supplied from the outside at the low heat treatment temperatures in FIGS. 94 and 95.

In the magnetic permeability (μ ') of FIG. 98, it turns out that the measured value which is 50000 in the heat processing temperature (Ta) about 500-600 degreeC, and the heating rate is about 20-200 degreeC / min is observed at injection temperature 1320 degreeC, and it turns out that it is favorable. At the injection temperature of 1280 ° C., a good value is also measured in a slightly narrower region as in FIG. 96.

The coercive force Hc in FIG. 99 is good because the heat treatment temperature Ta is approximately 510 to 540 ° C. at the injection temperature of 1320 ° C. and the temperature is about 40 m Oe at a temperature increase rate of 20 to 100 ° C./min. Although the coercive force (Hc) of 40 mOe or less is not measured at injection temperature of 1280 degreeC, it can be said that it is the same as the case of injection temperature of 1320 degreeC normally.

In the magnetic permeability µ 'in FIG. 100, the heat treatment temperature Ta is good at about 510 to 550 ° C. and the temperature increase rate at about 20 to 200 ° C./min. In addition, about the coercive force Hc in FIG. 101, about 45 mOe is measured in the place where heat processing temperature Ta intersects about 510-580 degreeC, and a temperature rising rate crosses about 40-200 degreeC / min.

In the magnetic permeability (μ ') of FIG. 102, the magnetic permeability (μ') of about 50000 is observed at the injection temperature of 1320 degreeC about 500-550 degreeC of heat processing temperature, and the temperature increase rate of about 20-40 degreeC / min. . Even at an injection temperature of 1280 ° C, a magnetic permeability of 500,000 units is observed.

The coercive force Hc in FIG. 103 is 40 to 49 m Oe at an injection temperature of 1320 ° C. at a heat treatment temperature Ta of about 500 to 600 ° C. and a temperature increase rate of about 20 to 200 ° C./min. The same coercive force Hc is observed in a slightly narrower region.

104 and 105 show the heat treatment temperature Ta and the heating rate, the permeability μ '(Fig. 104) and the coercive force (Hc) (Fig. 13) for the alloy having the composition Fe _85.5 Zr ₂ Nb ₄ B _8.5 described above. ) Is shown collectively for the holding time 0 to 60 min. Injection temperature is 1280 ℃. 106 and 107 are selected and shown only for the holding time 0, 10 and 60 min from FIGS. 104 and 105, and FIGS. 108 and 109 are the same and select only the holding time of 5 and 30 min. 110 to 115 show the same effect as those of Figs. 104 to 109 when the injection temperature is 1320 占 폚.

In these figures, the range considered to be particularly good in the magnetic permeability (μ ') is irrespective of the length and the length of the holding time. When the injection temperature is 1280 ° C, the heat treatment temperature Ta is around 500 to 560 ° C, and the temperature increase rate is 40. It can be seen that it is an intersection point of about 100 ° C./min (permeability (μ ′) ≒ 50000). In addition, when the injection temperature was 1320 ° C, it was also mentioned in the description of FIGS. 94 to 103, but it can be seen that it is wider than the good range when the injection temperature is 1280 ° C. In addition, in this case, it can be clearly seen that the permeability (μ ') exceeds about 30000 in almost all areas.

The range which can be said especially favorable about the coercive force Hc is the place where the heat processing temperature Ta intersects 520-560 degreeC and a temperature increase rate about 100 degree-C / min, when injection temperature is 1280 degreeC. ) ≒ 40 m Oe). Moreover, when injection temperature is 1320 degreeC, it turns out that it is wider than the range when injection temperature is 1280 degreeC similarly to permeability (micro '). In particular, it can be said that an appropriate range of expansion at the temperature increase rate is seen. Moreover, it turns out that the place where the coercive force Hc is about 50 mOe is distributed in a comparatively wide range.

From these figures, the influence of the holding time on the permeability (μ ') and the coercive force (Hc) of the holding time itself can be seen. On the other hand, it seems that the proper range tends to narrow as the holding time becomes longer. For example, when the injection temperature is 30 min or more at 1280 ° C, the area of 50000 is lost.

In view of the above, in order to obtain good permeability (μ ') and coercive force (Hc) regardless of the injection temperature and the holding time, the heat treatment temperature (Ta) is usually about 490 ° C or more and 670 ° C or less, more preferably. It is clear that the temperature may be set to 500 ° C or more and 560 ° C or less, and the temperature increase rate is usually 10 ° C / min or more and 200 ° C / min or less, more preferably 30 ° C / min or more and 100 ° C / min or less.

At this time, the following two advantages can be enjoyed from the lower limit of the temperature increase rate being 10 ° C / min or 30 ° C / min. That is, as a first point, being able to manufacture a highly saturated magnetic flux density low iron loss Fe-based soft magnetic alloy even if the temperature increase rate is as small as described above, it is possible to suppress equipment such as a furnace involved in heat treatment at a low cost. . Moreover, as a 2nd point, since the alloy in a heat processing process heats up slowly, it becomes possible to perform the heat processing without a spot as a whole. From these, it can be said that it is reasonable to select a condition in which the temperature rises as late as possible.

By the way, it turns out that the composition Fe _85.5 Zr ₂ Nb ₄ B _8.5 of the alloy in the figure demonstrated so far corresponds to the alloy in the first case. That is, since b = 85.5≥80 atomic% and x = 2, y = 4, x + y is 5 atomic% or more and 7.5 atomic% or less, and x / (x + y) is 1.5 / 6 or more and 2.5 / 6 or less Moreover, since it is 5 atomic% <=== 8.5 <= 12.5 atomic%, this satisfies all the conditions of a 1st alloy. In addition, the total amount of the composition ratio of Zr and the composition ratio of Nb in the alloy of the present composition is 6 atomic%, which is consistent with the above-mentioned "more preferable conditions (5.7 atomic% or more and 6.5 atomic% or less)" relating to the total amount. It is also.

116 and 117 show the relationship between the heat treatment temperature (Ta) and the temperature increase rate, the magnetic permeability (μ ') (FIG. 30) and the coercive force (Hc) (FIG. 31) in the alloy of composition Fe _85.5 Zr ₂ Nb ₄ B _8.5 . It is about the result of the measurement. In addition, a holding time is 5 minutes, and a thick solid line shows that an injection temperature is 1260 degreeC, and a thin solid line shows that an injection temperature is 1300 degreeC, respectively.

118 and 119 are drawings similar to the above, except that the composition of the alloy is (Fe _85.5 Zr ₂ Nb ₄ B _8.5 ) ₉₉ Zn ₁ . In addition, in the following FIGS. 120 and 121 and 122 and 123, the amount of Zn in the alloy composition is gradually increased, and in this order (Fe _85.5 Zr ₂ Nb ₄ B _8.5 ) ₉₈ Zn ₂ , ( Fe _85.5 Zr ₂ Nb ₄ B _8.5 ) ₉₇ Zn ₃ is a target alloy composition, and the magnetic permeability (μ ') and the coercive force (Hc) are shown.

However, (Fe _85.5 Zr ₂ Nb ₄ B _8.5 ) ₉₉ Zn ₁ , (Fe _85.5 Zr ₂ Nb ₄ B _8.5 ) ₉₈ Zn ₂ , and (Fe _85.5 Zr ₂ Nb ₄ B _8.5 ) ₉₇ Zn _{3 in} the alloy composition. Regarding the amount of Zn in the formula, this means that 1 atom%, 2 atom%, and 3 atom% of Zn were "injected" at the time of preparation of the alloy. That is, it does not mean that these alloys "contain" Zn as much as the above atomic%. This is because most of Zn evaporated and lost when producing a Zn-containing alloy as described above, and only a part of the injected amount remains in the alloy. And the amount of Zn actually contained in the said alloy is as showing in Table 8.

As shown in Table 8, when the amount of Zn injected in the case of Fe _85.5 Zr ₂ Nb ₄ B _8.5 is 1 atomic%, 0.056 atomic% Zn remains, and when 2 atomic% is 0.108 atomic% and 3 atoms as well. In the case of%, it can be seen that Zn of 0.133 to 0.141 atomic% remains, respectively. It goes without saying that these values correspond to "z" when the alloy in the third case is described. Further, these values naturally satisfy the above conditions, that is, 0.025 atomic% or more and 0.2 atomic% or less. In addition, although Table 8 shows the thing other than the composition related to FIGS. 32-37, these data may refer also later when explaining the various characteristics of the alloy containing Zn.

Then, first, the magnetic permeability (μ ') in FIG. 116 is measured in the range of 700-39800, and especially in the area | region in which the magnetic permeability (micro') becomes 35000 or more (injection temperature 1260 degreeC), heat processing temperature Ta is 510. About 550 degreeC and the place where the temperature increase rate is about 30-100 degreeC or about 200 degreeC / min cross | intersect. In addition, the coercive force Hc in FIG. 117 is measured in the range of 46-754 mOe, and the part which the heat processing temperature Ta cross | intersects 520 degreeC or more and a temperature increase rate of 50-100 degreeC or about 200 degreeC / min This can be said to be good.

In addition, the permeability (micro ') in FIG. 118 is measured in the range of 800-55700 (about 30000 or more), and it turns out that the highest value is larger than the previous FIG. In addition, about the coercive force Hc in FIG. 119, it is 37-670 mOe, and it turns out that the minimum value is smaller than the previous FIG. Thus, it can be seen that the soft magnetic properties were improved by adding Zn, and both the magnetic permeability (μ ') and the coercive force (Hc) were generally good.

In addition, the particularly preferable region in FIGS. 118 and 119 does not appear to be significantly different from FIGS. 116 and 117.

The magnetic permeability µ 'in FIG. 120 is slightly higher than that in the previous case when the maximum value of the measurement is 57500. In addition, it is understood that the permeability (μ ') is good in this case that the value of around 40000 or more is observed over the almost shown area | region. In addition, the coercive force Hc in FIG. 121 is measured in the range of 37-219 mOe, and the minimum value is slightly falling. However, since it is not so different from FIG. 119, it is favorable generally.

In addition, the magnetic permeability (μ ') in FIG. 122 is the highest value 52600, and the coercive force (Hc) in FIG. 123 is the minimum value 43 mOe, and these can also be said that it is not so different from the case mentioned above. In this case, however, the magnetic permeability (μ ') is slightly lower than the above-mentioned case, and the coercive force (Hc) is slightly higher than the above-mentioned case, and the overall soft magnetic properties appear to be slightly deteriorated.

The particularly good region in FIGS. 120 to 123 is not significantly different from that in FIGS. 118 and 119, that is, as long as it is irrelevant to the injection temperature.

124 and 125 show in one drawing that the injection amounts ζ of Zn are 0, 1, 2 and 3 when the injection temperature is 1260 ° C based on the data shown in FIGS. 116 to 123 described above. The relationship between the heat treatment temperature Ta and the temperature increase rate and the permeability μ '(Fig. 124) and the coercive force (Hc) (Fig. 125) is shown. 126 and 127 show the magnetic permeability mu '(FIG. 126) and the coercive force Hc (FIG. 127) at the injection temperature of 1300 DEG C for the same purpose as those of FIGS.

In these figures, the effect of adding Zn is easier to see compared to the previous one. That is, in FIG. 124, for example, an alloy having a composition containing no Zn exists in a region surrounding a region having a permeability (μ ') of 38000. In the case of ζ = 1, a region of 45000 is shown, and in FIG. 125, no Zn is contained. In contrast, the region of 48 m Oe was the lowest, whereas the region of 38 m Oe is shown at ζ = 1. The same can be said in FIG. 126 and FIG.

In general, it has been found that the soft magnetic properties are improved by adding Zn. The heat treatment temperature (Ta) and the temperature increase rate are not particularly different, and it is understood that an overall good alloy can be produced. That is, from this result, it can be said that it is preferable if it is about 510-550 degreeC as heat processing temperature Ta, and about 30-100 degreeC or about 200 degreeC / min as a temperature increase rate.

Example 7 Second Manufacturing Method

In the same manner as in Example 6, an alloy ribbon having a width of 15 mm and a thickness of about 20 μm was prepared. In addition, although the quenched ribbon is composed of an alloy mainly composed of amorphous materials, in order to precipitate fine crystal grains of bccFe and to improve soft magnetic properties, it is heated to a crystallization temperature or higher, and then subjected to a heat treatment to cool. A soft magnetic alloy ribbon sample of an iron-lossed Fe group and a sample of each comparative example are obtained.

The magnetic permeability (μ ') (real part of complex permeability) of the soft magnetic alloy ribbon sample obtained by the above method is processed into a ring shape having an outer diameter of 10 mm and an inner diameter of 6 mm by stacking the ribbon, and the windings are stacked and the impedance Measure with an analyzer. The measurement condition of the magnetic permeability (μ ') is 5 m Oe, 1 ㎑. In addition, the coercive force (Hc) and the saturation magnetic flux density (B ₁₀ ) are measured by measuring 10 Oe using a direct current BH loop tracer.

Table 9 and Table 10 show the (200) planes of the bccFe phase by X-ray diffraction for the alloy before the heat treatment (the alloy immediately after solidification on the cooling roll or after the quench or quenched state) for each of the ribbon samples obtained through the above steps. The results of the measurement of strength, coercive force (Hc [Oe]), saturation magnetic flux density (B ₁₀ [T]), residual magnetic flux density (Br [T]), permeability (μ ') (real part of complex permeability) (The contents of Ta and T _x1 , T _x2 and T _{x1 '} will be described later). In addition, "start" in Table 10 indicates that the alloy was formed immediately after the ribbon began to be formed on the cooling roll, that is, the ejection of the melt from the nozzle was started. Indicates an alloy.

As can be seen from these tables, it can be seen that the composition of the highly saturated magnetic flux density low iron loss Fe-based soft magnetic alloy shown here corresponds to the first case. That is, it can be seen that each of Fe ₈₅ Zr _1.75 Nb _4.25 B ₉ , Fe _85.5 Zr ₂ Nb ₄ B _8.5 , and Fe ₈₅ Zr ₂ Nb _4.5 B _8.5 , and all satisfy the conditions for the first alloy. In the following, these will be examined focusing on the contrast between the injection temperature of less than 1350 ° C and that of 1350 ° C or more.

First, in the alloy before the heat treatment of Fe ₈₅ Zr _1.75 Nb _4.25 B ₉ , the X-ray diffraction was performed on both the start side and the end side with respect to the injection temperature below 1350 ° C. As a result, no peak due to the crystal phase was detected. On the other hand, the peak from the (200) plane of bccFe phase was observed about what became injection temperature 1350 degreeC or more. In Table 1, the former is represented by "-" and the latter by "XX". Incidentally, "X" means that the peak was not as strong as that represented by "XX".

As such, when the injection temperature is lower than 1350 ° C., the alloy to be produced is almost completely amorphous, and in the opposite case, at least a part of the crystallization is already made, despite being before heat treatment. Since it is important for the present invention that the alloy structure before the heat treatment is amorphous, the above result is one proof that an alloy having a more preferable injection temperature of less than 1350 ° C. can be produced.

For the coercive force (Hc), 0.038 to 0.044 (starting side) for the injection temperature below 1350 ° C, 0.038 to 0.044 (end side), 0.76 to 0.086 (starting side), 0.049 for the injection temperature above 1350 ° C. To 0.078 (end) and permeability (μ ') at 40968 to 49672 (start), 41508 to 49649 (end) at injection temperatures below 1350 ° C, 23812 to 24739 (starting at injection temperatures above 1350 ° C). ) And 25594 to 38191 (end side), it can be seen that an alloy having a preferable injection temperature of less than 1350 ° C is produced. In addition, since it is confirmed that the saturation magnetic flux density (B ₁₀ ) in the case where these injection temperatures are less than 1350 ° C. is 1.5 T or more in almost all regardless of the start side and the end side, it can be said to have excellent soft magnetic properties in that respect.

In addition, it can be seen in Table 9 and Table 10 that the Fe _85.5 Zr ₂ Nb ₄ B _8.5 , Fe ₈₅ Zr ₂ Nb _4.5 B _8.5 is almost the same as above. In other words, by lowering the injection temperature to 1350 ° C., an alloy exhibiting excellent soft magnetic properties with small coercive force, high permeability, and sufficient saturation magnetic flux density can be produced.

By the way, the T _x1 [℃] is the crystallization temperature of the bccFe in Table 9, T _x2 [℃] represents the crystallization temperature of the compound to be crystallized in a high temperature side of the T _x1, respectively. In addition, T _{x1 '[℃]} is a crystallization temperature of the other compound that takes the median value of the T _x1 and T _x2. These T _x1 , T _x2 , and T _{x1 ′} are factors that determine the parameters relating to heat treatment as described above. That is, as described above,? Tx = T _x2- T _x1 = T _x1- T _{x1 '} ?

It can be seen that the highly saturated magnetic flux density low iron loss Fe-based soft magnetic metal of each composition in Table 9 and Table 10 satisfies ΔTx? In other words, it can be seen that the heat treatment in these alloys is appropriately made.

Next, Table 11, Table 12, Table 13, and Table 14 are demonstrated. Since each alloy described in these tables does not contain Zn, it can be corresponded to the alloy in a said 1st case. In these tables, the injection temperature is lower than 1350 ° C in all cases.

First, the coercive force (Hc) can be seen in the range of 0.038 to 0.116 (Oe: start side) and 0.043 to 0.114 (Oe: end) when looking at Tables 11, 12, 13, and 14 as a whole. In most cases, the value is in the range of 0.04 Oe. Subsequently, the saturation magnetic flux density (B ₁₀ ) is 1.5 T or more in almost all cases. (The ribbon Lots RQ 6-108 and RQ 6-111 in Table 11 and Table 12 are those that are 1.5 T or more.) End, only the start side of the same RQ 6-148 in Table 13 and Table 14). In addition, there are some exceptions to the magnetic permeability μ ', but most of them have values exceeding 30000.

As a result, it can be seen that the highly saturated magnetic flux density low iron loss Fe-based soft magnetic alloys prepared at the injection temperatures of 1350 ° C. are excellent in the soft magnetic properties. In this case as well, the heat treatment temperature Ta, such as the crystallization temperatures T _x1 and T _x2 , is, of course, to satisfy the above conditions described in Table 9.

By the way, it is considered that these soft magnetic properties are particularly excellent in Tables 9 to 14, for example, Fe ₈₅ Zr _1.75 Nb _4.25 B ₉ , Fe _85.5 Zr ₂ Nb ₄ B _8.5 , Table 11 and Tables 9 and 10. Fe _85.25 Zr _1.75 Nb _4.25 B _8.75 , Fe _85.75 Zr _2.25 Nb _3.75 B _8.25 of Table 12, Fe ₈₆ Zr _2.25 Nb _3.75 B ₈ of Table 13 and Table 14, Fe _85.62 Zr ₂ Nb ₄ B _8.38, etc. are mentioned. In these cases, it turns out that the total amount of the composition ratio of Zr and the composition ratio of Nb is 6 atomic%.

Next, Table 15 and Table 16 are demonstrated. These Tables 15 and 16 show the high saturation magnetic flux density low iron loss Fe-based soft magnetic alloys having a composition containing Zn, that is, the ribbon samples of the alloys as in the second to third cases. In the same manner as described in Fig. 6, the results of the measurement of various properties on matters such as coercive force (Hc), saturation magnetic flux density (B ₁₀ ), and permeability (μ ') are shown.

By the way, in the amounts of Zn in the composition formulas shown in Tables 15 and 16, "1", "2", and "3" described in the subscript are 1 atomic%, 2 atomic%, And 3 atomic% of Zn are "inserted". That is, it does not mean that these alloys "contain" Zn for each atomic%. As described above, since most of Zn evaporated and lost when producing an alloy containing Zn, only a part of the injected amount remains as an alloy. And the amount of Zn actually contained in the said alloy is as showing in Table 8.

As can be seen from Tables 15 and 16, the injection temperature is less than 1350 ° C in all cases. The result is that the coercive force (Hc) is 0.037 to 0.131 (start), 0.040 to 0.140 (end), the saturation magnetic flux density (B ₁₀ ) is 1.59 or more (start), 1.52 or more (end), and the permeability (μ ') 23088 to 63337 (start side) and 20936 to 55250 (end side). As for the magnetic permeability (μ '), as described now, the 20000 units also appear sporadically, but in total, more than 30,000 units or 40000 to 50000 units occupy most of them. Therefore, it can be seen from Tables 15 and 16 that both highly saturated magnetic flux density low iron loss Fe-based soft magnetic alloys having good soft magnetic properties are manufactured.

In addition, no peak due to the (200) plane is observed in all cases except for the ribbon Lot RQ 6-172 and RQ 6-170. Therefore, it is suggested that most of the samples posted here exhibit a good structure mainly composed of an amorphous phase.

In addition, it turns out that the said conditions regarding crystallization temperature ( _Tx1 , _Tx2 ), heat processing temperature (Ta), etc. are satisfied also in the case of these Tables 15 and 16.

Example 8: Third Manufacturing Method

In the same manner as in Example 1, after producing a ribbon in a quenched state having a composition of Fe _85.5 Zr ₂ Nb ₄ B _8.5 and Fe ₇₈ Si ₉ B ₁₃ (commercially available: amorphous alloy), the first heat treatment was performed. On the other hand, in the same manner as in Example 5 (Fe _0.8575 Zr _0.02 Nb _0.04 B _0.0825 ) _99.88 Zn _0.12 , (Fe _0.8575 Zr _0.02 Nb _0.04 B _0.0825 ) _99.87 Zn _0.13 after producing a ribbon of the quenched state composition After the heat treatment, the temperature is lowered to room temperature.

Here, the first heat treatment condition of the ribbon is Fe _85.5 Zr ₂ Nb ₄ B _8.5 , (Fe _0.8575 Zr _0.02 Nb _0.04 B _0.0825 ) _99.88 Zn _0.12 , (Fe _0.8575 Zr _0.02 Nb _0.04 B _0.0825 ) _99.87 Zn _0.13 180 degree-C / min, holding time 5 minutes, and holding temperature are performed on the conditions of 510 degreeC, 525 degreeC, and 510 degreeC, respectively. The Fe ₇₈ Si ₉ B ₁₃ was carried out under the conditions of a heating rate of 180 ° C./min, a holding time of 120 minutes and a holding temperature of 350 ° C.

Next, each obtained ribbon is subjected to a second heat treatment in air, and at this time, the coercive force and permeability of each sample when the holding time at the holding temperature of 320 ° C. is changed to a range of 0 to 100 hours, and B ₁₀ and Br ( Residual magnetization) is measured. Further, the second heat treatment time dependency of these magnetic properties is investigated. The second heat treatment time dependence of the magnetic properties here examines the second heat treatment time (holding time), the coercive force and permeability, and the rate of change of B ₁₀ and Br (residual magnetic flux density). B ₁₀ here is substantially the same as the saturation magnetic flux density. In addition, the 2nd heat processing conditions of a ribbon here are 20 degree-C / min of the temperature increase rate at the time of temperature rising from the room temperature to 320 degreeC which is a holding temperature. The results are shown in Tables 17-20. In addition, FIG. 128 shows the second heat treatment time dependency of the coercive force Hc. 129 shows the second heat treatment time dependence of the magnetic permeability. In Tables 17-20 and FIGS. 128-129, the change rate of each magnetic characteristic is a change rate with respect to this based on the magnetic characteristic when a 2nd heat treatment time is 0 hours, ie, not performing a 2nd heat treatment. It represents.

As is clear from the results shown in Table 17 and FIG. 128, the second heat treatment was performed after the first heat treatment of the ribbon in a quenched state having a composition of Fe ₇₈ Si ₉ B ₁₃ (commercially available: amorphous alloy). As the coercive force exceeds 0.05 Oe and the second heat treatment time becomes longer, the coercive force changes significantly. In contrast, the ribbon in the quenched state of Fe _85.5 Zr ₂ Nb ₄ B _8.5 in the composition range of the present invention, (Fe _0.8575 Zr _0.02 Nb _0.04 B _0.0825 ) _99.88 Zn ribbon in the quenched state of _0.12 , (Fe _0.8575 Zr _0.02 Nb _0.04 B _0.0825 ) _99.87 Zn After _performing the first heat treatment on the ribbons having the quenched composition of _0.13 , the coercive force of the second heat treatment for 1 hour or more is smaller than that of the non-coating. . As a result, in the sample within the composition range of the present invention and the production method of the present invention, the coercive force decreases, whereas in the comparative example sample having the composition of Fe ₇₈ Si ₉ B ₁₃ , the coercive force is significantly deteriorated. Able to know. As a result of this, it can be seen that the application of the production method of the present invention is an alloy having a fine crystal structure, preferably an alloy within the composition range of the present invention.

As apparent from the results shown in Table 18 and FIG. 129, the second heat treatment was performed after the first heat treatment of the ribbon in a quenched state of Fe ₇₈ Si ₉ B ₁₃ (commercially available product: amorphous alloy). It is 5200 or less, and permeability changes drastically as the 2nd heat treatment time becomes long. In contrast, the ribbon in the quenched state of Fe _85.5 Zr ₂ Nb ₄ B _8.5 in the composition range of the present invention, (Fe _0.8575 Zr _0.02 Nb _0.04 B _0.0825 ) _99.88 Zn ribbon in the quenched state of _0.12 , (Fe _0.8575 Zr _0.02 Nb _0.04 B _0.0825 ) _99.87 Zn The second heat treatment after the first heat treatment of the ribbon of the composition of the composition of _0.13 after the first heat treatment, the permeability even if the permeability is greater than 38500 and the second heat treatment time is longer It can be seen that the rate of change of is small. As a result, when considered together with the results of Table 17 and FIG. 128, the sample within the composition range of the present invention and the manufacturing method of the present invention has a small coercive force, an improved permeability, or maintains a high value. Therefore, it can be seen that the soft magnetic characteristics are improved. On the other hand, the comparative example of the composition of Fe ₇₈ Si ₉ B ₁₃ , the permeability is greatly degraded when the production method of the present invention is applied, the production method of the present invention is an alloy having a fine crystal structure, preferably, It is preferably applied to alloys within the composition range of the invention.

As is apparent in the results shown in Table 19 Fe ₇₈ Si ₉ B ₁₃ (commercially available product: amorphous alloy), after a quench condition ribbon of the composition subjected to the first heat treatment, is subjected to the second heat treatment, the B ₁₀ is 1.57 or less It is small and the change rate of B ₁₀ is small. In contrast, the ribbon in the quenched state of Fe _85.5 Zr ₂ Nb ₄ B _8.5 in the composition range of the present invention, (Fe _0.8575 Zr _0.02 Nb _0.04 B _0.0825 ) _99.88 Zn ribbon in the quenched state of _0.12 , (Fe _0.8575 Zr _0.02 Nb _0.04 B _0.0825 ) _99.87 Zn After the first heat treatment of the ribbon in the quenched state composition of _0.13 , the second heat treatment is performed, even if B ₁₀ is larger than 1.57 and the second heat treatment time is long. You can see this little thing. It can be seen from this that the sample within the composition range of the present invention and subjected to the third production method of the present invention maintains a much higher saturation magnetic flux density than the comparative example sample having the composition of Fe ₇₈ Si ₉ B ₁₃ . .

As apparent from the results shown in Table 20, the second heat treatment time was performed after the first heat treatment of the ribbon in a quenched state of Fe ₇₈ Si ₉ B ₁₃ (commercially available amorphous alloy). As the length increases, the residual magnetic flux density changes drastically. In contrast, the ribbon in the quenched state of Fe _85.5 Zr ₂ Nb ₄ B _8.5 in the composition range of the present invention, (Fe _0.8575 Zr _0.02 Nb _0.04 B _0.0825 ) _99.88 Zn ribbon in the quenched state of _0.12 , (Fe _0.8575 Zr _0.02 Nb _0.04 B _0.0825 ) _99.87 Zn After the first heat treatment of the ribbons of the quenched composition of _0.13 , the second heat treatment is performed, and the change rate of the residual magnetic flux density is small even if the second heat treatment time is longer. have. Thereby, the sample which is in the composition range of this invention and performed the 3rd manufacturing method of this invention can improve soft magnetic property and raise residual magnetic flux density. On the other hand, the comparative example sample of the composition of Fe ₇₈ Si ₉ B ₁₃ can increase the residual magnetic flux density by applying the manufacturing method of the present invention, but it can be seen that the soft magnetic properties deteriorate significantly.

From the results of FIGS. 128 to 129 and Tables 17 to 20, when the production method of the present invention is applied to an alloy having a fine crystal structure, preferably, an alloy shown in the composition range of the present invention, high permeability, saturation magnetic flux density, It is possible to reduce the coercive force while maintaining the residual magnetic flux density.

At the same time, it is also found that by applying the manufacturing method of the present invention, a soft magnetic alloy having a small change over time can be provided as long as the heat is lower than the second heat treatment temperature.

As a result of the present invention as described above, it is possible to provide a high saturation magnetic flux density low iron loss Fe-based soft magnetic alloy or magnetic core excellent in soft magnetic properties, small iron loss, large breakage deformation, workability.

In addition, the iron loss can be 0.10 w / kg (at 1.4 T, 50 kPa) or less, and the saturation magnetic flux density can be 1.5 T or more. It can cope with the bending process at the time of manufacture, etc., and can provide the high saturation magnetic flux density low iron loss Fe base soft magnetic alloy or magnetic core which is suitable for a trans use etc.

In addition, it was possible to provide a method for producing an alloy having excellent low coercive force, high permeability, and high saturation magnetic flux density characteristics, to develop a method for producing a soft magnetic alloy according to the patent application, to improve soft magnetic properties, and to maintain a long time at high temperature. Even if left unattended, there is little effect of change in the magnetic properties over time, and it is possible to provide a method for producing a Fe-based soft magnetic alloy suitable for a transformer or the like that can cope with processing during the production of a transformer or the like.

Claims (58)

An alloy mainly composed of an amorphous phase is heat-treated, and at least 50% of the structure is a microcrystalline structure mainly composed of a bcc-Fe phase having an average grain size of 100 nm or less, and a highly saturated magnetic flux density low iron loss Fe having a composition shown in the following formula. Key Magnetic Alloys. Fe _a Zr _x Nb _y B _z Where a, x, y, z indicating the composition ratio is 80 atomic% ≤ a, 5 atomic% ≤ x + y ≤ 7 atomic%, 1.5 / 6 ≤ x / (x + y) ≤ 2.5 / 6, 5 atoms % ≤z≤12.5 atomic%. The method according to claim 1, wherein a, x, y, z representing the composition ratio is 83 atomic% ≤ a, 5.7 atomic% ≤ x + y ≤ 6.5 atomic%, 1.5 / 6 ≤ x / (x + y) ≤ 2.5 / A high saturated magnetic flux density low iron loss Fe-based soft magnetic alloy, characterized in that 6, 6 atomic% ≤z≤9.5 atomic%. The crystallization temperature of the bcc-Fe phase is T _x1 , the crystallization temperature of the compound phase crystallized at a higher temperature side than T _x1 is T _x2 , and the interval (ΔT _x ) of the crystallization temperatures is ΔT _x = T. _A highly saturated magnetic flux density low iron loss Fe-based soft magnetic alloy having a temperature of 200 ° C. ΔΔT _x when _{x 2} −T _{x 1} . 2. The saturation magnetic flux density according to claim 1, wherein the saturation magnetic flux density is 1.5 T or more, and the iron loss when the magnetic flux of 1.4T is applied at a frequency of 50 Hz is 0.15 W / kg or less, and the iron loss change rate before and after aging at 200 ° C. for 500 hours is High saturation magnetic flux density low iron loss Fe base soft magnetic alloy of 10% or less. The high saturation magnetic flux density low iron loss Fe-based soft magnetic alloy according to claim 1, wherein the fracture strain is 1.0 × 10 ⁻² or more. Comprising a composition represented by the following formula, at least 50% or more of the structure is made of microcrystalline grains of bccFe having an average grain size of 100nm or less, the remainder is made of amorphous alloy phase, the microcrystalline grains of bccFe to quench the alloy, A high saturation magnetic flux density low iron loss Fe-based magnetic alloy made of a nearly amorphous single-phase structure and then precipitated after cooling the amorphous phase above a crystallization temperature. (Fe _1-a Q _a ) _b B _x M _y Zn _z Provided that Q is either Co, Ni or both, and M is one or two or more elements selected from Ti, Zr, Hf, V, Nb, Ta, Mo, and W; , x, y, z is a = 0, 80 atomic% ≤b, 5 atomic% ≤x≤12.5 atomic%, 5 atomic% ≤y≤7 atomic%, 0.025 atomic% ≤z≤0.2 atomic%. The high saturation magnetic flux density low iron loss Fe-based soft magnetic alloy according to Claim 6, wherein the iron loss change rate at 320 ° C. for 100 hours is 20% or less, the saturation magnetic flux density is 1.5T or more, and the magnetic permeability is 30000 or more. 7. The highly saturated magnetic flux density low iron loss Fe-based soft magnetic alloy according to claim 6, wherein the fracture strain is 10 × 10 ⁻³ or more. Comprising a composition represented by the following formula, at least 50% or more of the structure is made of microcrystalline grains of bccFe having an average grain size of 100nm or less, the remainder is made of amorphous alloy phase, the microcrystalline grains of bccFe to quench the alloy, A high saturation magnetic flux density low iron loss Fe-based magnetic alloy obtained by forming an amorphous single phase structure and cooling the amorphous phase after heating above a crystallization temperature. (Fe _1-a Q _a ) _b B _x M _y Zn _z M ' _u Provided that Q is either Co, Ni, or both, M is one or two or more elements selected from Ti, Zr, Hf, V, Nb, Ta, Mo, and W, and M 'is Cr, Ru. And one or two or more elements selected from R h and Ir, and a, b, x, y, z, and u representing the composition ratio are a = 0, 80 atomic% ≤ b, 5 atomic% ≤ x ≤ 12.5 atomic%, 5 atomic% ≤ y ≤ 7 atomic%, 0.025 atomic% ≤ z ≤ 0.2 atomic%, u ≤ 5 atomic%. The high saturation magnetic flux density low iron loss Fe-based soft magnetic alloy according to claim 9, wherein a change rate of iron loss by heating at 320 ° C. for 100 hours is 20% or less, a saturation magnetic flux density of 1.5T or more, and a permeability of 30000 or more. The high saturation magnetic flux density low iron loss Fe-based soft magnetic alloy according to claim 9, wherein the fracture strain is 10 × 10 ⁻³ or more. Comprising a composition represented by the following formula, at least 50% or more of the structure is made of microcrystalline grains of bccFe having an average grain size of 100nm or less, the remainder is made of amorphous alloy phase, the microcrystalline grains of bccFe to quench the alloy, A high saturation magnetic flux density low iron loss Fe-based magnetic alloy, wherein the amorphous phase is formed into a single-phase structure, and the amorphous phase is cooled after being heated to a crystallization temperature or higher. Fe _b Zr _x Nb _y B _t Zn _z However, b, x, y, t, and z representing the composition ratio are 80 atomic% ≤ b, 1.5 atomic% ≤ x ≤ 2.5 atomic%, 3.5 atomic% ≤ y ≤ 5.0 atomic%, 5 atomic% ≤ t ≤ 12.5 atoms %, 0.025 atomic% ≤ z ≤ 0.2 atomic%, 5.0 atomic% ≤ x + y ≤ 7.5 atomic%, and 1.5 / 6 ≤ x / (x + y) ≤ 2.5 / 6. The high saturation magnetic flux density low iron loss Fe-based soft magnetic alloy according to claim 12, wherein a change rate of iron loss by heating at 320 ° C. for 100 hours is 20% or less, a saturation magnetic flux density of 1.5T or more, and a permeability of 30000 or more. The method according to claim 12, wherein the composition is represented by the following formula, wherein at least 50% or more of the tissue is composed of microcrystalline grains of bccFe having an average grain size of 100 nm or less, the balance consisting of amorphous alloy phase, the microcrystalline grains of the bccFe A high saturation magnetic flux density low iron loss Fe-based soft magnetic alloy in which a silver alloy is quenched to form a nearly amorphous single-phase structure, and the amorphous phase is cooled after being heated above a crystallization temperature. (Fe _1-a Q _a ) _b Zr _x Nb _y B _t Zn _z However, Q is either Co, Ni, or both, and a, b, x, y, t, and z, which represent the composition ratio, a ≤ 0.05, 80 atomic% ≤ b, 1.5 atomic% ≤ x ≤ 2.5 atomic%, 3.5 atomic% ≤ y ≤ 5.0 atomic%, 5 atomic% ≤ t ≤ 12.5 atomic%, 0.025 atomic% ≤ z ≤ 0.2 atomic% , 5.0 atomic% ≤ x + y ≤ 7.5 atomic% and 1.5 / 6 ≤ x / (x + y) ≤ 2.5 / 6. The method according to claim 12, wherein the composition is represented by the following formula, wherein at least 50% or more of the tissue is composed of microcrystalline grains of bccFe having an average grain size of 100 nm or less, the balance consisting of amorphous alloy phase, the microcrystalline grains of the bccFe A high saturation magnetic flux density low iron loss Fe-based soft magnetic alloy in which a silver alloy is quenched to form an almost amorphous single-phase structure, and the amorphous phase is cooled after being heated above a crystallization temperature. (Fe _1-a Q _a ) _b Zr _x Nb _y B _t Zn _z M ' _u Provided that Q is either Co, Ni, or both, and M 'is one, two or more elements selected from Cr, Ru, Rh, and Ir, and represents a, b, x, y, t, and z, indicating a composition ratio. , u is, a ≤ 0.05, 80 atomic% ≤ b, 1.5 atomic% ≤ x ≤ 2.5 atomic%, 3.5 atomic% ≤ y ≤ 5.0 atomic%, 5 atomic% ≤ t ≤ 12.5 atomic%, 0.025 atomic% ≤ z ≤ 0.2 atomic%, u ≦ 5 atomic%, 5.0 atomic% ≦ x + y ≦ 7.5 atomic%, 1.5 / 6 ≦ x / (x + y) ≦ 2.5 / 6. 13. The highly saturated magnetic flux density low iron loss Fe-based soft magnetic alloy according to claim 12, wherein the fracture strain is 1.0 × 10 ⁻² or more. An alloy mainly composed of an amorphous phase is heat-treated, so that at least 50% of the structure is a microcrystalline structure mainly composed of a bcc-Fe phase having an average grain size of 100 nm or less, and a low iron loss Fe soft magnetic composition having a composition represented by the following formula: Low iron loss core made of alloy. Fe _a Zr _x Nb _y B _z Where a, x, y, z indicating the composition ratio is 80 atomic% ≤ a, 5 atomic% ≤ x + y ≤ 7.5 atomic%, 1.5 / 6 ≤ x / (x + y) ≤ 2.5 / 6, 5 atoms % ≤z≤12.5 atomic%. 18. The method according to claim 17, wherein a, x, y, z representing the composition ratio of the low iron loss Fe-based soft magnetic alloy is 83 atomic% ≤ a, 5.7 atomic% ≤ x + y ≤ 6.5 atomic%, 1.5 / 6 ≤ x Low iron loss magnetic core, characterized in that / (x + y) ≤ 2.5 / 6, 6 atomic% ≤ z ≤ 9.5 atomic%. The method of claim 17, wherein the interval of the low iron loss Fe-based soft crystallization temperature of the compound to the bcc-Fe crystallization temperature of the magnetic alloy with T _x1, and the crystallization at a high temperature side of the T _x1 to T _x2 and crystallization temperature When (ΔT _x ) is ΔT _x = T _x2 −T _x1 , the low iron loss magnetic core is 200 ° C. ≦ ΔT _x . 18. The saturation magnetic flux density according to claim 17, wherein the saturation magnetic flux density is 1.5 T or more, the iron loss when the magnetic flux of 1.4T is applied at a frequency of 50 Hz is 0.15 W / kg or less, and the iron loss change rate before and after aging at 200 ° C. for 500 hours is 10. Low iron loss magnetic core using low iron loss Fe-based soft magnetic alloy of less than%. 18. The low iron loss magnetic core according to claim 17, wherein the fracture strain is made of a ribbon of a low iron loss Fe-based soft magnetic alloy of 1.0 × 10 ⁻² or more. 18. The low iron loss magnetic core according to claim 17, wherein one or two or more annular bodies formed of a ribbon of said low iron loss Fe-based soft magnetic alloy are laminated. 18. The low iron loss magnetic core according to claim 17, wherein the low iron loss magnetic core is made of an annular shape wound around a ribbon of the Fe-based soft magnetic alloy. With a composition represented by the following formula, at least 50% or more of the structure is composed of microcrystalline grains of bccFe having an average grain size of 100 nm or less, the balance consisting of amorphous alloy phases, and the microcrystalline grains of bccFe quench the alloy, thereby almost amorphous phase. The low iron loss magnetic core consisting of a low iron loss Fe-based magnetic alloy, characterized in that after the amorphous phase of the amorphous phase is cooled after heating above the crystallization temperature. (Fe _1-a Q _a ) _b B _x M _y Zn _z Provided that Q is either Co, Ni, or both, and M is one, two or more elements selected from Ti, Zr, Hf, V, Nb, Ta, Mo, and W; x, y, z is a = 0, 80 atomic% ≤b, 5 atomic% ≤x≤12.5 atomic%, 5 atomic% ≤y≤7.5 atomic%, 0.025 atomic% ≤z≤0.2 atomic%. 25. The low iron loss as set forth in claim 24, wherein the iron loss at the time of heating at 200 ° C. for 500 hours is 10% or less, when the magnetic flux of 1.4T is applied at a saturation magnetic flux density of 1.5T or more and a frequency of 50 Hz is 0.15 W / kg or less. Low iron loss core using Fe-based soft magnetic alloy. 25. The low iron loss magnetic core according to claim 24, wherein the fracture strain of the low iron loss Fe-based soft magnetic alloy is 1.0 × 10 ⁻² or more. 25. The low iron loss magnetic core according to claim 24, wherein one or two or more annular bodies formed of said low iron loss Fe-based soft magnetic ribbon are laminated. 25. The low iron loss magnetic core according to claim 24, wherein the low iron loss magnetic core comprises a ring wound around the ribbon of the low iron loss Fe-based soft magnetic alloy to form an annular shape. With a composition represented by the following formula, at least 50% or more of the structure is composed of microcrystalline grains of bccFe having an average grain size of 100 nm or less, the balance consisting of amorphous alloy phases, and the microcrystalline grains of bccFe quench the alloy, thereby almost amorphous phase. A low iron loss magnetic core consisting of a low iron loss Fe-based soft magnetic alloy which is formed by heating the amorphous phase after heating to a crystallization temperature or higher after forming into a single phase structure of. (Fe _1-a Q _a ) _b B _x M _y Zn _z M ' _u Provided that Q is either Co, Ni or both, M is one or more elements selected from Ti, Zr, Hf, V, Nb, Ta, Mo, W, and M 'is Cr, Ru And one or two or more elements selected from R h and Ir, and a, b, x, y, z, and u representing the composition ratio are a = 0, 80 atomic% ≤ b, 5 atomic% ≤ x ≤ 12.5 atomic%, 5 atomic% ≤ y ≤ 7.5 atomic%, 0.025 atomic% ≤ z ≤ 0.2 atomic%, u ≤ 5 atomic%. 30. The iron loss according to claim 29, wherein, in the low iron loss magnetic core, a rate of change of iron loss by heating at 200 ° C. for 500 hours is 10% or less, a saturation magnetic flux density of 1.5T or more, and a frequency of 1.4T at a frequency of 50 Hz is applied. Low iron loss magnetic core using low iron loss Fe-based soft magnetic alloy of 0.15 W / kg or less. 30. The low iron loss magnetic core according to claim 29, wherein the low iron loss Fe-based soft magnetic alloy has a fracture strain of 1.0 × 10 ⁻² or more. The low iron loss magnetic core according to claim 29, wherein one or two or more annular bodies formed of said low iron loss Fe-based soft magnetic ribbon are laminated. 30. The low iron loss magnetic core according to claim 29, wherein the low iron loss magnetic core is made of a ring wound around a ribbon of the Fe-based soft magnetic alloy. With a composition represented by the following formula, at least 50% or more of the structure consists of microcrystalline grains of bccFe having an average grain size of 100 nm or less, the balance consisting of amorphous alloy phases, wherein the microcrystalline grains of bccFe quench the alloy, and are almost amorphous. The low iron loss magnetic core consisting of a low iron loss Fe-based soft magnetic alloy which is precipitated after being heated to a crystallization temperature or higher after being formed into a single-phase structure of. Fe _b Zr _c Nb _d B _x Zn _z However, b, c, d, x, and z representing the composition ratio is 80 atomic% ≤ b, 1.5 atomic% ≤ c ≤ 2.5 atomic%, 3.5 atomic% ≤ d ≤ 5.0 atomic%, 5 atomic% ≤ x ≤ 12.5 atoms %, 0.025 atomic% ≤ z ≤ 0.2 atomic%, 5.0 atomic% ≤ c + d ≤ 7.5 atomic%, and 1.5 / 6 ≤ c / (c + d) ≤ 2.5 / 6. 35. The composition according to claim 34, wherein at least 50% or more of the structure consists of microcrystal grains of bccFe having an average grain size of 100 nm or less, the balance consisting of amorphous alloy phases, and the microcrystal grains of bccFe are alloys. The low iron loss magnetic core consisting of a low iron loss Fe-based magnetic alloy, which is quenched and formed into a single-phase structure of almost amorphous phase, and then cooled after the amorphous phase is heated above the crystallization temperature. (Fe _1-a Q _a ) _b Zr _c Nb _d B _x Zn _z However, Q is one of Co, Ni, or both, and a, b, c, d, x, and z representing the composition ratio is a≤0.05, 80 atomic% ≤b, 1.5 atomic% ≤c≤2.5 atomic%, 3.5 atomic% ≤d≤5.0 atomic%, 5 atomic% ≤x≤12.5 atomic%, 0.025 atomic% ≤z≤0.2 atomic% , 5.0 atomic% ≦ c + d ≦ 7.5 atomic% and 1.5 / 6 ≦ c / (c + d) ≦ 2.5 / 6. The method of claim 34, wherein the composition represented by the following formula, wherein at least 50% or more of the tissue is composed of microcrystalline grains of bccFe having an average grain size of 100nm or less, the remainder is made of amorphous alloy phase, the microcrystalline grains of bccFe, A low iron loss magnetic core made of a low iron loss Fe-based magnetic alloy, wherein the alloy is quenched and formed into an almost amorphous single phase structure, and then the amorphous phase is cooled after being heated above the crystallization temperature. (Fe _1-a Q _a ) _b Zr _c Nb _d B _x Zn _z M ' _u Provided that Q is either Co, Ni, or both, and M 'is one or two or more elements selected from Cr, Ru, Rh, and Ir, and represents a, b, c, d, x, z , u is, a≤0.05, 80 atomic% ≤b, 1.5 atomic% ≤c≤2.5 atomic%, 3.5 atomic% ≤d≤5.0 atomic%, 5 atomic% ≤x≤12.5 atomic%, 0.025 atomic% ≤z≤0.2 atomic%, u ≦ 5 atomic%, 5.0 atomic% ≦ c + d ≦ 7.5 atomic%, 1.5 / 6 ≦ c / (c + d) ≦ 2.5 / 6. 35. The low iron loss as set forth in claim 34, wherein the iron loss when a change in iron loss by heating at 200 ° C. for 500 hours is 10% or less, a saturation magnetic flux density of 1.5T or more, and a 1.4T magnetic flux is applied at a frequency of 50 Hz is 0.15 W / kg or less. Low iron loss core using Fe-based soft magnetic alloy. 35. The low iron loss magnetic core according to claim 34, wherein the fracture strain of the low iron loss Fe-based soft magnetic alloy is 1.0 × 10 ⁻² or more. The low iron loss magnetic core according to claim 34, wherein one or two or more annular bodies formed of said low iron loss Fe-based soft magnetic ribbon are laminated. 35. The low iron loss magnetic core according to claim 34, wherein the low iron loss magnetic core is formed of a ring wound around a ribbon of the Fe-based soft magnetic alloy. delete delete delete Melt melted in a crucible with an alloy containing one or two or more elements M and B selected from Ti, Zr, Hf, V, Nb, Ta, Mo, W, and Mn in a crucible containing Fe as a main component. Injected from the nozzle to the cooling roll at a temperature of quenching, quenching and solidifying on the cooling roll to form a ribbon mainly composed of amorphous, heat treatment at a temperature above the crystallization temperature, at least 50% or more of the structure is bccFe having an average grain size of 100 nm or less A method for producing a Fe-based magnetic alloy composed of fine grains mainly composed of microstructures having a structure in which the balance is amorphous. 45. The method of producing a Fe-based soft magnetic alloy according to claim 44, wherein the melt injection temperature is at least 1240 占 폚. The average grain size of the amorphous alloy containing Fe as the main component and one or two or more elements M and B selected from Ti, Zr, Hf, V, Nb, Ta, Mo, W, and Mn is subjected to the first heat treatment. Fe group having a fine bcc structure of 30 nm or less, mainly made of a microcrystalline alloy containing an amorphous phase, and then performing a second heat treatment at a holding temperature of 100 ° C. or higher and a holding temperature of the first heat treatment temperature or lower. Method for producing soft magnetic alloy. 47. The method for producing an Fe-based soft magnetic alloy according to claim 46, wherein the holding temperature of the second heat treatment is 200 to 400 deg. 47. The method for producing an Fe-based magnetic alloy according to claim 46, wherein the second heat treatment is performed for 0.5 to 100 hours. 49. The method for producing an Fe-based soft magnetic alloy according to claim 48, wherein the second heat treatment is maintained for 1 to 30 hours. 47. The method of claim 46, wherein the first heat treatment is performed at a temperature increase rate of 10 to 200 deg. C / min. 47. The method for producing an Fe-based soft magnetic alloy according to claim 46, wherein the holding temperature of said first heat treatment is 500 to 800 deg. Comprising a composition represented by the following formula, at least 50% or more of the structure is made of microcrystalline grains of bccFe having an average grain size of 100nm or less, the remainder is made of amorphous alloy phase, the microcrystalline grains of bccFe to quench the alloy, A high saturation magnetic flux density low iron loss Fe-based magnetic alloy made of a nearly amorphous single-phase structure and then precipitated after cooling the amorphous phase above a crystallization temperature. (Fe _1-a Q _a ) _b B _x M _y Zn _z Provided that Q is either Co, Ni or both, and M is one or two or more elements selected from Ti, Zr, Hf, V, Nb, Ta, Mo, and W; , x, y, z is 0 <a≤0.05, 80 atomic% ≤b, 5 atomic% ≤x≤12.5 atomic%, 5 atomic% ≤y≤7 atomic%, 0.025 atomic% ≤z≤0.2 atomic%. Comprising a composition represented by the following formula, at least 50% or more of the structure is made of microcrystalline grains of bccFe having an average grain size of 100nm or less, the remainder is made of amorphous alloy phase, the microcrystalline grains of bccFe to quench the alloy, A high saturation magnetic flux density low iron loss Fe-based magnetic alloy obtained by forming an amorphous single phase structure and cooling the amorphous phase after heating above a crystallization temperature. (Fe _1-a Q _a ) _b B _x M _y Zn _z M ' _u Provided that Q is either Co, Ni, or both, M is one or two or more elements selected from Ti, Zr, Hf, V, Nb, Ta, Mo, and W, and M 'is Cr, Ru. And one or two or more elements selected from R h and Ir, and a, b, x, y, z, and u representing the composition ratio are 0 <a≤0.05, 80 atomic% ≤b, 5 atomic% ≤x≤12.5 atomic%, 5 atomic% ≤y≤7 atomic%, 0.025 atomic% ≤z≤0.2 atomic%, u≤5 atomic%. The method according to claim 12, wherein the composition is represented by the following formula, wherein at least 50% or more of the tissue is composed of microcrystalline grains of bccFe having an average grain size of 100 nm or less, the balance consisting of amorphous alloy phase, the microcrystalline grains of the bccFe A high saturation magnetic flux density low iron loss Fe-based soft magnetic alloy in which a silver alloy is quenched to form a nearly amorphous single-phase structure, and the amorphous phase is cooled after being heated above a crystallization temperature. (Fe _1-a Q _a ) _b Zr _x Nb _y B _t Zn _z However, Q is either Co, Ni, or both, and a, b, x, y, t, and z, which represent the composition ratio, 0 <a≤0.05, 80 atomic% ≤b, 1.5 atomic% ≤x≤2.5 atomic%, 3.5 atomic% ≤y≤5.0 atomic%, 5 atomic% ≤t≤12.5 atomic%, 0.025 atomic% ≤z≤0.2 atomic %, 5.0 atomic% ≤ x + y ≤ 7.5 atomic%, and 1.5 / 6 ≤ x / (x + y) ≤ 2.5 / 6. With a composition represented by the following formula, at least 50% or more of the structure is composed of microcrystalline grains of bccFe having an average grain size of 100 nm or less, the balance consisting of amorphous alloy phases, and the microcrystalline grains of bccFe quench the alloy, thereby almost amorphous phase. The low iron loss magnetic core consisting of a low iron loss Fe-based magnetic alloy, characterized in that after the amorphous phase of the amorphous phase is cooled after heating above the crystallization temperature. (Fe _1-a Q _a ) _b B _x M _y Zn _z Provided that Q is either Co, Ni, or both, and M is one, two or more elements selected from Ti, Zr, Hf, V, Nb, Ta, Mo, and W; x, y, z is 0 <a≤0.05, 80 atomic% ≤b, 5 atomic% ≤x≤12.5 atomic%, 5 atomic% ≤y≤7.5 atomic%, 0.025 atomic% ≤z≤0.2 atomic%. With a composition represented by the following formula, at least 50% or more of the structure is composed of microcrystalline grains of bccFe having an average grain size of 100 nm or less, the balance consisting of amorphous alloy phases, and the microcrystalline grains of bccFe quench the alloy, thereby almost amorphous phase. A low iron loss magnetic core consisting of a low iron loss Fe-based soft magnetic alloy which is formed by heating the amorphous phase after heating to a crystallization temperature or higher after forming into a single phase structure of. (Fe _1-a Q _a ) _b B _x M _y Zn _z M ' _u Provided that Q is either Co, Ni or both, M is one or more elements selected from Ti, Zr, Hf, V, Nb, Ta, Mo, W, and M 'is Cr, Ru And one or two or more elements selected from R h and Ir, and a, b, x, y, z, and u representing the composition ratio are 0 <a ≦ 0.05, 80 atomic% ≦ b, 5 atomic% ≦ x ≦ 12.5 atomic%, 5 atomic% ≦ y ≦ 7.5 atomic%, 0.025 atomic% ≦ z ≦ 0.2 atomic%, u ≦ 5 atomic%. 35. The composition according to claim 34, wherein at least 50% or more of the structure consists of microcrystal grains of bccFe having an average grain size of 100 nm or less, the balance consisting of amorphous alloy phases, and the microcrystal grains of bccFe are alloys. The low iron loss magnetic core consisting of a low iron loss Fe-based magnetic alloy, which is quenched and formed into a single-phase structure of almost amorphous phase, and then cooled after the amorphous phase is heated above the crystallization temperature. (Fe _1-a Q _a ) _b Zr _c Nb _d B _x Zn _z However, Q is one of Co, Ni, or both, and a, b, c, d, x, and z representing the composition ratio is 0 <a≤0.05, 80 atomic% ≤b, 1.5 atomic% ≤c≤2.5 atomic%, 3.5 atomic% ≤d≤5.0 atomic%, 5 atomic% ≤x≤12.5 atomic%, 0.025 atomic% ≤z≤0.2 atomic %, 5.0 atomic% ≤ c + d ≤ 7.5 atomic%, and 1.5 / 6 ≤ c / (c + d) ≤ 2.5 / 6. The method of claim 34, wherein the composition represented by the following formula, wherein at least 50% or more of the tissue is composed of microcrystalline grains of bccFe having an average grain size of 100nm or less, the remainder is made of amorphous alloy phase, the microcrystalline grains of bccFe, A low iron loss magnetic core made of a low iron loss Fe-based magnetic alloy, wherein the alloy is quenched and formed into an almost amorphous single phase structure, and then the amorphous phase is cooled after being heated above the crystallization temperature. (Fe _1-a Q _a ) _b Zr _c Nb _d B _x Zn _z M ' _u Provided that Q is either Co, Ni, or both, and M 'is one or two or more elements selected from Cr, Ru, Rh, and Ir, and represents a, b, c, d, x, z , u is, 0 <a≤0.05, 80 atomic% ≤b, 1.5 atomic% ≤c≤2.5 atomic%, 3.5 atomic% ≤d≤5.0 atomic%, 5 atomic% ≤x≤12.5 atomic%, 0.025 atomic% ≤z≤0.2 atomic %, u ≦ 5 atomic%, 5.0 atomic% ≦ c + d ≦ 7.5 atomic%, and 1.5 / 6 ≦ c / (c + d) ≦ 2.5 / 6.