JP4974603B2 - Rotor core of motor and manufacturing method thereof - Google Patents

Description

The present invention relates to a rotor core of a motor having high yield strength and excellent magnetic characteristics, and a method for manufacturing the same.

As represented by main motors of electric vehicles and hybrid electric vehicles, recent motors are strongly required to save energy and improve efficiency.
Increasing the frequency of the motor is one effective means for achieving energy savings and higher efficiency. However, as the frequency increases, the rotational speed of the motor also increases. Since the centrifugal force applied to the core constituting the rotor increases as the motor speed increases, the core material is required to have a high yield strength. In other words, when the yield strength of the core material is insufficient, the core material undergoes plastic deformation due to centrifugal force, and the air gap between the rotor core and the stator core changes from the design value, and the motor performance deteriorates. During rotation, the rotor and the stator come into contact with each other, resulting in damage to the motor. For this reason, it is indispensable to increase the strength of the rotor core material in order to save energy and increase the efficiency of the motor by increasing the frequency.

  By the way, as a method of manufacturing a rotor core, a method of laminating electromagnetic steel sheets having a thickness of about 0.35 to 0.5 mm has been generally used so far. Since it takes a great deal of money to stack several hundred sheets, it has begun to consider putting a motor into practical use that uses electromagnetic steel bars that do not require stacking instead of electromagnetic steel sheets.

  However, the current electromagnetic bar steel is made of low carbon steel or silicon steel like the electromagnetic steel sheet, and the strength is not necessarily high because the main strengthening mechanism is the solid solution strengthening of ferrite. For example, even in the case of 3 mass% Si steel, the yield strength is about 350 MPa. Moreover, although it is an example of a steel plate, as disclosed in Patent Document 1, the yield strength is approximately 300 to 450 MPa even for the purpose of increasing the strength, and further, after performing cold forging, annealing is performed. Even if it gives, sufficient yield strength is not obtained.

Conventionally, carbon steel for mechanical structures and chromium molybdenum steel used as cold forged parts can obtain high yield strength depending on the heat treatment conditions when quenching and tempering treatment is performed after cold forging, but the magnetic properties are low. The problem was that it was very low.
JP 2002-371340 A

  Accordingly, an object of the present invention is to provide a cold forged part having a sufficient magnetic property as a rotor core material and having a high yield strength, and a manufacturing method thereof.

The inventors have made the structure of the cold forged part a ferrite single phase, and further dispersed and precipitated fine precipitates with a particle size of less than 10 nm in the ferrite structure to strengthen the precipitation of ferrite, thereby significantly increasing the strength and magnetic properties. It has been found that a rotor core of a motor having excellent characteristics can be obtained.

This invention is made | formed based on the said knowledge, The summary structure is as follows.
1. C: 0.040 mass% or more and 0.120 mass% or less,
Si: 0.5% by mass or less,
Mn: 0.60 mass% or more and 3.00 mass% or less,
Al: 0.1 mass% or less,
Ti: 0.03 mass% to 0.35 mass% and
Mo: 0.05 mass% or more and 0.8 mass% or less, having a component composition of the balance Fe and inevitable impurities, having a structure in which fine precipitates having a particle size of less than 10 nm are dispersed in ferrite, {211} A rotor core for a motor, comprising a cold forged part having a strength ratio of {100} face to the face of 0.7 or more and a strength ratio of {110} face to the {211} face of 0.4 or less.

2. 2. The rotor core of the motor according to 1 above, wherein the component composition satisfies the following formula (1).
Record
0.50 ≦ (C / 12) / [(Ti / 48) + (Mo / 96)] ≦ 1.50 (1)

3. 3. The rotor core of the motor according to 1 or 2, wherein the fine precipitate is a carbide of Ti and / or Mo.

4). As the component composition,
Nb: 0.08 mass% or less,
4. The rotor core of the motor according to 1, 2 or 3, characterized in that it comprises one or more selected from V: 0.15 mass% or less and W: 1.5 mass% or less.

5. 5. The rotor core of the motor according to 4 above, wherein the component composition satisfies the following formula (2).
Record
0.50 ≦ (C / 12) / [(Ti / 48) + (Mo / 96) + (Nb / 93) + (V / 51) + (W / 184)] ≦ 1.50 (2)

6). 6. The rotor core of the motor according to 4 or 5, wherein the fine precipitate is a carbide containing Ti and Mo and at least one of Nb, V and W.

7). The component composition further includes S: 0.01% by mass or more and 0.1% by mass or less, and
Pb: 0.2 mass% or less,
Ca: 0.005 mass% or less,
The rotor core of the motor according to any one of 1 to 6 above, which contains one or more of Bi: 0.1% by mass or less and B: 0.02% by mass or less.

8). C: 0.040 mass% or more and 0.120 mass% or less,
Si: 0.5% by mass or less,
Mn: 0.60 mass% or more and 3.00 mass% or less,
Al: 0.1 mass% or less,
Ti: 0.03 mass% to 0.35 mass% and
Mo: A steel material containing 0.05% by mass to 0.8% by mass with the balance being Fe and inevitable impurities is heated to 1100 ° C or higher and hot-rolled at a finishing temperature of 880 ° C or higher. subjecting% or more cold forging, and facilities annealed at 600 ° C. or higher 700 ° C. or less of the temperature range, {211} {100} intensity ratio of surface 0.7 or more with respect to surface and {211} with respect to plane {110} A method for manufacturing a rotor core of a motor, wherein the strength ratio of the surface is 0.4 or less .

9. 9. The method for manufacturing a rotor core of a motor according to 8, wherein the steel material has a component composition satisfying the following formula (1).
Record
0.50 ≦ (C / 12) / [(Ti / 48) + (Mo / 96)] ≦ 1.50 (1)

Ten. The steel material is further
Nb: 0.08 mass% or less,
10. The method for producing a rotor core of a motor as described in 8 or 9 above, comprising one or more selected from V: 0.15% by mass or less and W: 1.5% by mass or less.

11． 11. The method for producing a rotor core of a motor as described in 10 above, wherein the steel material has a composition that satisfies the following formula (2).
Record
0.50 ≦ (C / 12) / [(Ti / 48) + (Mo / 96) + (Nb / 93) + (V / 51) + (W / 184)] ≦ 1.50 (2)

12． The steel material further includes S: 0.01 mass% or more and 0.1 mass% or less, and
Pb: 0.2 mass% or less,
Ca: 0.005 mass% or less,
12. The method for producing a rotor core of a motor as described in any one of 8 to 11 above, wherein one or more of Bi: 0.1% by mass or less and B: 0.02% by mass or less are included.

  According to the present invention, a cold forged component having sufficient magnetic properties as a rotor core material and high yield strength can be provided. Therefore, in a motor using the cold forged component of the present invention as a rotor core, the rotation of the motor Even if the speed is increased, the above-mentioned problems caused by plastic deformation of the core material can be avoided. Therefore, the frequency in the motor can be further increased, and the energy saving and high efficiency of the motor are realized. Therefore, the present invention is extremely useful industrially.

Hereinafter, the cold forged part of the present invention will be described in detail below for each component composition, microstructure and production condition. In addition, unless otherwise indicated, the "%" display regarding a component composition means "mass%".
C: 0.040% or more and 0.120% or less When C is less than 0.040%, the amount of fine precipitates is insufficient and high yield strength cannot be obtained, so C needs to be 0.040% or more. On the other hand, if the content of C exceeds 0.120%, the precipitates become coarse and high yield strength cannot be obtained, so the upper limit of C needs to be 0.120%.

Si: 0.5% or less
Since Si decreases the cold workability, the addition amount is limited to 0.5% or less. More preferably, it is 0.15% or less.

Mn: 0.60% or more and 3.00% or less In the present invention, the precipitation behavior of precipitates is closely related to the progress of transformation from austenite to ferrite (hereinafter referred to as ferrite transformation). When the difference between the transformation start temperature and the precipitation start temperature of the precipitate is small and the ferrite transformation and precipitation compete with each other, the precipitate is finely dispersed and precipitated in the ferrite. Mn contributes to competing ferrite transformation and precipitation by lowering the ferrite transformation temperature and reducing the difference between the transformation initiation temperature of the ferrite transformation and the precipitation initiation temperature of the precipitate. Furthermore, Mn contributes to high strength by solid solution strengthening, and this effect becomes remarkable when Mn is added in excess of 0.60%. On the other hand, when the amount of Mn exceeds 3.00%, a low-temperature transformation phase such as bainite is generated in addition to ferrite, and strengthening by fine precipitates is insufficient, and the magnetic flux density is lowered. For this reason, the upper limit of Mn is set to 3.00%. More preferably, it is 0.70% or more and 2.80% or less.

Al: 0.1% or less
Al may be added as a deoxidizing element. In this case, 0.010% or more needs to be added. However, if it is added excessively, not only the effect is saturated, but the amount of AlN that is a precipitate with N increases, and AlN does not precipitate below 10 nm. In order to avoid this, the amount of Al added is 0.1% or less. More preferably, it is 0.05% or less.

Ti: 0.03% to 0.35%
Ti is an effective component for improving the strength by finely depositing precipitates including Ti-based carbides and Ti-Mo-based carbides, and 0.03% or more is necessary to ensure high yield strength. On the other hand, if added over 0.35%, the precipitates become coarse and the strength decreases, so Ti is made 0.03% to 0.35%. More preferably, it is 0.03% or more and 0.20% or less.

Mo: 0.05% or more and 0.8% or less
Mo is added to finely precipitate precipitates including Mo-based carbides and Ti-Mo-based carbides, and improve strength. Further, Mo has a slow diffusion rate, and when it precipitates together with Ti, it has an advantage that the growth rate of the precipitate is reduced and a fine precipitate is easily obtained. Here, in order to ensure high yield strength, it is necessary to add 0.05% or more of Mo. On the other hand, when adding over 0.8%, a low-temperature transformation phase such as bainite is generated in addition to ferrite. Precipitation strengthening due to fine precipitates is insufficient, resulting in a decrease in strength and a deterioration in magnetic properties. For this reason, Mo is made 0.05% to 0.8%. More preferably, it is 0.15% or more and 0.50% or less.

Further, in the above component composition, particularly regarding the atomic ratio of the amounts of C, Ti and Mo, satisfying the following formula (1) is advantageous for making the precipitate finer.
Record
0.50 ≦ (C / 12) / [(Ti / 48) + (Mo / 96)] ≦ 1.50 (1)
This parameter affects the size of the precipitate, and when it is set to 0.50 or more and 1.50 or less, formation of fine precipitates having a particle diameter of less than 10 nm is facilitated, which is preferable.

  It was observed that Ti and Mo in the carbides of fine Ti—Mo based carbides have an atomic ratio of Ti / Mo of 0.2 or more and 2.0 or less, and finer carbides of 0.7 or more and 1.5 or less.

As described above, although the essential components have been described, in the present invention, in order to further improve the strength and toughness,
One or more of Nb, V and W can be added.
Nb: 0.08% or less
Nb forms fine precipitates together with Ti and Mo and contributes to an increase in strength. Moreover, ductility and toughness are improved by adjusting the grain size of ferrite. For that purpose, 0.01% or more is preferably added. However, if the content exceeds 0.08%, the ferrite becomes finer and fine precipitates adversely affect the magnetic properties, so the addition amount should be 0.08% or less. More preferably, it is 0.04% or less.

V: 0.15% or less V is also added in an amount of 0.01% or more because V also forms fine precipitates with Ti and Mo and contributes to an increase in strength. However, if the content exceeds 0.15%, the precipitate becomes coarse, so the addition amount is 0.15% or less. More preferably, it is 0.10% or less.

W: 1.5% or less Since W also forms fine precipitates together with Ti and Mo and contributes to an increase in strength, W is preferably added at 0.1% or more. However, if the content exceeds 1.5%, the precipitate becomes coarse, so the addition amount is 1.5% or less. More preferably, it is 1.0% or less.

When these elements are added, if the atomic ratios of these elements and the amounts of C, Ti, and Mo are defined according to the following formula (2), it is advantageous for refinement of precipitates.
Record
0.50 ≦ (C / 12) / [(Ti / 48) + (Mo / 96) + (Nb / 93) + (V / 51) + (W / 184)] ≦ 1.50 (2)

  This parameter affects the size of the precipitates. When the parameter is 0.50 or more and 1.50 or less, formation of fine precipitates having a particle size of less than 10 nm is facilitated.

  In addition, in the fine carbide | carbonized_material containing 1 type, or 2 or more types of Nb, V, and W, atomic ratio (Ti + Nb + V) / (Mo + W) of Ti, Mo, Nb, V, and W in carbide is 0.2-2.0, and more delicate It was observed to be 0.7-1.5 for carbides.

Further, in the present invention, in order to improve the machinability at the time of processing the part, S is included in the range of 0.01% to 0.1%, Pb: 0.2% or less, Ca: 0.005% or less, Bi: 0.1% or less, and B : One or more of 0.02% or less can be added.
Here, the reason why the S content is 0.01% or more and 0.1% or less is that if the S content is less than 0.01%, the machinability cannot be improved, whereas if it exceeds 0.1%, the ductility and toughness deteriorate. Because.
S is less than 0.01% and is contained as an impurity. In the present invention, strength and magnetic properties are not affected at a content of 0.1% or less. Therefore, it can add actively and it can be set as 0.01 to 0.1% of content.

  Also, for Pb, Ca, Bi, and B, when the addition amount exceeds the respective upper limit, ductility and toughness are reduced. Therefore, the addition amount is Pb ≦ 0.2%, Ca ≦ 0.005%, Bi ≦ 0.1% and It is necessary to set B ≦ 0.02%. The lower limit is preferably set to Pb ≧ 0.05%, Ca ≧ 0.001%, Bi ≧ 0.03%, and B ≧ 0.001% for reasons of improving machinability.

  In addition, for the purpose of improving ductility and toughness, one or more of Cr, Ni and Cu may be added in the range of Cr ≦ 0.5%, Ni ≦ 0.5% and Cu ≦ 0.5%, respectively. If added, both should be 0.05% or more.

That is, P and N, which are unavoidable impurities, are elements that are not preferable for the magnetic properties, so it is desirable to reduce P and N as much as possible. Specifically, it is preferable to restrict P to 0.03% or less. N is preferably regulated to 0.01% or less, more preferably 0.005% or less.
In addition, the effect of this invention is not impaired by the presence or absence and content of these elements.

  Next, in the present invention, the microstructure is defined as a ferrite single phase. The reason why the ferrite single phase is used is that the ferrite single phase is the most preferable structure for magnetic properties. The ferrite single phase in the present invention means that the area ratio of ferrite is 95% or more, preferably 98% or more in cross-sectional structure observation (200-times optical microscope structure observation).

  Furthermore, in the present invention, the particle size of the fine precipitate is less than 10 nm. When the grain size of the precipitate is 10 nm or more, the precipitation strengthening ability is insufficient and high yield strength cannot be obtained. In other words, the smaller the particle size of the fine precipitates, the more effective the strength is, and it is desirably 5 nm or less, more desirably 3 nm or less. As such fine precipitates, carbides containing a composite of Ti and Mo or carbides further containing one or more of Nb, V and W are preferable. The fine precipitates are precipitated during cooling after hot rolling and during subsequent annealing.

As for the number of fine precipitates, 1000 / μm ³ or more, more desirably 5000 / μm ³ or more is preferable because high yield strength is easily obtained.
Although the distribution form of these fine precipitates is not particularly defined, it is desirable to uniformly disperse and precipitate in the matrix.

  In addition to the precipitates described above, the effect of the present invention is not impaired even if a small amount of Fe carbide is contained. However, if a large amount of Fe carbide having an average particle size of 1 μm or more is contained, the magnetic properties are hindered. In this case, the upper limit of the size of the Fe carbide contained is preferably 1 μm, and the content is preferably 1% or less of the entire precipitate.

The size of the precipitate and the ratio of the fine precipitate to the total precipitate are obtained by the following method.
That is, an electron microscope sample is prepared by an electrolytic polishing method using a twin jet method, and the sample is observed at an acceleration voltage of 200 kV. At that time, the focus is shifted from the normal focus in order to control the crystal orientation of the parent phase so that the precipitate has a measurable contrast with respect to the parent phase and to minimize the number of precipitates. Observe by method. Moreover, the sample thickness of the area | region which measured the deposit particle | grains is evaluated by measuring an elastic scattering peak and an inelastic scattering peak intensity using an electron energy loss spectroscopy.

By this method, the measurement of the particle diameter and the number of particles and the measurement of the sample thickness can be executed for the same region. The measurement of the particle diameter and the number of particles is carried out on four places of a 0.5 μm × 0.5 μm region of the sample, and the precipitates distributed per 1 μm ² are calculated as the number for each particle diameter. Next, from this value and the thickness of the sample, the number of precipitates per particle diameter distributed per 1 μm ³ is calculated. Thereby, the size of the precipitate and the proportion of the precipitate having a particle size of less than 10 nm in the total precipitate are obtained.

In the present invention, it is important that the structure has a strength ratio of {100} plane to {211} plane in the rolling direction of 0.7 or more and a strength ratio of {110} plane to {211} plane of 0.4 or less. is there. This is because the {100} plane includes many <100> easy axes that are most easily magnetized, while the {110} plane includes many <110> axes that are not easily magnetized. Therefore, as the strengthening ratio of the {100} plane increases, the spontaneous magnetization of the magnetic domain increases, and the magnetic flux density increases with a constant excitation current.
In order to quantify the influence of {100} orientation and {110} orientation on magnetic properties, the inventors analyzed the intensity ratio of {100} plane to {211} plane and {110} plane to {211} plane. As a result of investigating the magnetic flux density B ₅₀ for the samples with various strength ratios, it was found that the former had a high B ₅₀ value of 0.7 or more and the latter 0.4 or less (see FIG. 1). Therefore, in the present invention, the strength ratio of the {100} plane to the {211} plane is 0.7 or more, and the strength ratio of the {110} plane to the {211} plane is 0.4 or less.

Below, the conditions at the time of manufacturing the above-mentioned cold forged part are demonstrated. In other words, the cold forged parts were heated to 1100 ° C or higher after the steel material adjusted to the above-mentioned components, hot-rolled at a finishing temperature of 880 ° C or higher, and then subjected to cold forging with a processing rate of 20% or higher. It can be obtained by annealing in a temperature range of from ℃ to 700 ℃.
Heating temperature In the present invention, in order to precipitate precipitates finely during cooling after hot rolling, it is necessary to once dissolve the precipitates precipitated on the slab before hot rolling in a heating furnace. is there. At that time, if the heating temperature is less than 1100 ° C., Ti—Mo-based carbides and the like are not sufficiently dissolved, so the heating temperature is set to 1100 ° C. or higher.

Finishing temperature In the present invention, the precipitation behavior of the precipitate is closely related to the progress of the ferrite transformation, and the difference between the transformation start temperature of the ferrite transformation that occurs during cooling after rolling and the precipitation start temperature of the precipitate is small, When the ferrite transformation and precipitation compete, the precipitate is finely dispersed and precipitated in the ferrite. In order to compete with ferrite transformation and precipitation, it is necessary to lower the start temperature of ferrite transformation, but when the finishing temperature in hot rolling is low, the strain introduced by rolling raises the start temperature of ferrite transformation and precipitates. Inhibits refinement of objects. In order to avoid this, the finishing temperature may be set to a high temperature at which the influence of distortion does not appear. From this point, the finishing temperature is set to 880 ° C. or higher.

In the present invention, the strength ratio of the {100} plane to the {211} plane needs to be 0.7 or more and the strength ratio of the {110} plane to the {211} plane needs to be 0.4 or less. For this reason, it is important to develop a texture that is superior in magnetic properties by cold working and heat treatment following hot rolling, and at the same time sufficiently precipitate fine precipitates to combine high strength and good magnetic properties. It is.

Below, the experiment conducted in order to determine processing conditions is explained in full detail.
That is, the component composition is within the scope of the present invention, C: 0.072%, Si: 0.07%, Mn: 2.05%, Ti: 0.22%, Mo: 0.48%, P: 0.007%, S: 0.005%, Al: 0.018% And N: 0.0020%, and the balance was made of iron and steel with inevitable impurities. This steel was heated to 1230 ° C. and hot-rolled to obtain a steel bar having a length of 6 m and a diameter of 70 mm. Thereafter, a disk-shaped cold forging test piece having a diameter of 60 mm and a height of 25 mm was cut out so that the upsetting direction was perpendicular to the rolling direction, and upsetting was performed with the degree of processing varied from 0 to 80%. Annealing was carried out at 60 ° C. for 60 min.

The test pieces after cold forging and annealing thus obtained were measured for surface strength ratio, magnetic properties and mechanical properties.
First, the surface strength ratio was obtained by measuring 25 mm × 25 mm and a thickness of 1 mm from the center of the 1/2 position of the plate thickness, and measuring by X-ray using Cu—Kα ray. Regarding the magnetic characteristics, a ring-shaped test piece having an inner diameter of 33 mm, an outer diameter of 45 mm, and a thickness of 3 mm was taken from the center of the plate thickness, subjected to 100 primary windings and 100 secondary windings, and a DC excitation current of 5000 A. The magnetic flux density B ₅₀ at / m was measured. The mechanical properties were evaluated by a tensile test.

FIG. 1 shows the influence of the annealing temperature on the magnetic measurement results. From FIG. 1, in the steel of the present invention having a cold work rate of 20% or more, B ₅₀ was 1.66 T or more and YS was 650 MPa or more.
If the cold work rate is less than 20%, the development of {100} texture that is advantageous for the magnetic properties is insufficient, and the magnetic properties are also inferior because the {110} texture is high. Moreover, since precipitation of fine precipitates by cold forging and annealing is not promoted, the yield strength is not sufficiently improved. As a result of the above examination, the cold working rate was set to 20% or more.

Annealing temperature When the annealing temperature is 600 ° C or lower, the precipitation strength cannot be improved because fine precipitates do not precipitate. On the other hand, when the annealing temperature exceeds 700 ° C., the magnetic properties are drastically lowered due to coarsening of fine precipitates. For this reason, the annealing temperature is limited to 600 ° C. or more and 700 ° C. or less.

  Steels having the composition shown in Table 1 were melted, heated to 1200 ° C., and then hot-rolled to form steel bars having a length of 6 m and a diameter of 80 mm. At that time, the cooling rate after hot rolling was 0.23 ° C./s. Thereafter, a disk-shaped cold forging test piece having a diameter of 60 mm and a height of 25 mm was cut out so that the upsetting direction was perpendicular to the rolling direction, and upsetting was performed with the degree of processing varied from 0 to 80%. Annealing was performed at 650 ° C. for 60 minutes.

  About the test piece after the cold forging and annealing thus obtained, the surface strength ratio, the magnetic property and the mechanical property were measured in the same manner as described above.

Furthermore, a thin film sample was prepared by electrolytic polishing, and the particle size of the precipitate was measured by observation with a transmission electron microscope (TEM) according to the method described above, and an energy dispersive X-ray spectrometer (EDX) was used in combination. The precipitate was identified.
Table 2 shows the results of the above-described structure observation, tensile test, and magnetic measurement.

  In the structure section of Table 1, ferrite is abbreviated as T when a low temperature transformation phase such as F, bainite, martensite, etc. is generated and its volume fraction is 5% or more. Moreover, about the deposit, the average particle diameter was described. The variation in particle diameter was ± 1 nm at the maximum for precipitates of less than 10 nm, and ± 3 nm to ± 5 nm for precipitates larger than that. In the case where a low temperature transformation phase was generated in the structure, measurement of the crystal grain size and the particle size of the precipitate was omitted.

Table 2 shows the influence of the steel composition. As is clear from the table, the invention example yielded a yield strength of 600 MPa or more, and the magnetic properties were such that the magnetic flux density B ₅₀ was 1.67 T. Excellent with the above.

On the other hand, in the comparative example in which the steel composition is out of the scope of the present invention, the yield strength is low and the magnetic properties are inferior.
That is, No. 13 has a low C, an insufficient amount of fine precipitates, and a low yield strength.
No. 14 has high C, precipitates are coarsened, and yield strength is low. When the precipitate is coarse, the magnetic property is inferior because the precipitate adversely affects the magnetic property as described above.
In No. 15-A, since Mn is low, ferrite transformation and precipitation do not compete sufficiently, and as a result of the coarse precipitation, the strength is low and the magnetic properties are also deteriorated.
In No. 16-A having a high Mn, a low-temperature transformation phase is formed, and the yield strength is low because precipitation strengthening due to fine precipitates is insufficient. Also, the magnetic properties are low due to the generation of the low temperature transformation phase.
No. 17-A has a low yield strength due to a low amount of fine precipitates because Ti is low. On the other hand, in No. 18-A where Ti is high, precipitates are coarsened, yield strength is low, and magnetic flux density is low.
No. 19-A has a low Mo, so the amount of fine precipitates is insufficient and the yield strength is low. On the other hand, in No. 20-A having a high Mo, a low-temperature transformation phase is generated, and precipitation strengthening due to fine precipitates is insufficient, resulting in low yield strength and poor magnetic properties. The magnetic flux density is low as in No. 16-A, which has a high Mn and also generates a low-temperature transformation phase.

Next, Table 3 shows the results of manufacturing the steel No. 5, which is the steel of the present invention, under various conditions. As is clear from the table, the invention example has obtained a high yield strength of 650 MPa or more, As for the magnetic characteristics, the magnetic flux density B ₅₀ is an excellent value of 1.67 T or more.
On the other hand, No. 5-D steel has a high annealing temperature, so that the precipitate is dissolved and the second phase is precipitated during cooling. As a result, only low magnetic properties are shown.
Since No. 5-E steel is not annealed, precipitation due to annealing does not occur, and the strength is lower than that of the annealed material. Also, {100} texture does not develop and magnetic properties are low.
Since No. 5-F steel and N0.5-G steel have a low cold work rate, {100} texture does not develop sufficiently even after annealing, and magnetic properties are also low.
Since the heating temperature of No. 5-M steel is low, the slab precipitate before hot rolling does not sufficiently dissolve in the heating furnace, and the precipitate becomes coarse. As a result, the magnetic properties are inferior.
In No. 5-Q steel, the finishing temperature is low, and the strain introduced by rolling increases the starting temperature of the ferrite transformation and inhibits the competition between ferrite transformation and precipitation. As a result, the precipitates become coarse and the yield strength decreases. In addition, magnetic properties are deteriorated due to coarse precipitates.

It is a figure which shows the relationship between a cold work rate, a texture, a magnetic characteristic, and a mechanical characteristic.

Claims (12)

C: 0.040 mass% or more and 0.120 mass% or less,
Si: 0.5% by mass or less,
Mn: 0.60 mass% or more and 3.00 mass% or less,
Al: 0.1 mass% or less,
Ti: 0.03 mass% to 0.35 mass% and
Mo: 0.05 mass% or more and 0.8 mass% or less, having a component composition of the balance Fe and inevitable impurities, having a structure in which fine precipitates having a particle size of less than 10 nm are dispersed in ferrite, {211} A rotor core for a motor, comprising a cold forged part having a strength ratio of {100} face to the face of 0.7 or more and a strength ratio of {110} face to the {211} face of 0.4 or less. The rotor core of the motor according to claim 1, wherein the component composition satisfies the following expression (1).
Record
0.50 ≦ (C / 12) / [(Ti / 48) + (Mo / 96)] ≦ 1.50 (1) The rotor core of the motor according to claim 1, wherein the fine precipitate is a carbide of Ti and / or Mo. As the component composition,
Nb: 0.08 mass% or less,
4. The rotor core of a motor according to claim 1 , wherein the rotor core includes one or more selected from V: 0.15 mass% or less and W: 1.5 mass% or less. The rotor core of the motor according to claim 4, wherein the component composition satisfies the following expression (2).
Record
0.50 ≦ (C / 12) / [(Ti / 48) + (Mo / 96) + (Nb / 93) + (V / 51) + (W / 184)] ≦ 1.50 (2) The rotor core of a motor according to claim 4 or 5, wherein the fine precipitate is a carbide containing Ti and Mo and at least one of Nb, V and W. The component composition further includes S: 0.01% by mass or more and 0.1% by mass or less, and
Pb: 0.2 mass% or less,
Ca: 0.005 mass% or less,
The rotor core of the motor according to any one of claims 1 to 6, comprising one or more of Bi: 0.1 mass% or less and B: 0.02 mass% or less. C: 0.040 mass% or more and 0.120 mass% or less,
Si: 0.5% by mass or less,
Mn: 0.60 mass% or more and 3.00 mass% or less,
Al: 0.1 mass% or less,
Ti: 0.03 mass% to 0.35 mass% and
Mo: A steel material containing 0.05% by mass to 0.8% by mass with the balance being Fe and inevitable impurities is heated to 1100 ° C or higher and hot-rolled at a finishing temperature of 880 ° C or higher. subjecting% or more cold forging, and facilities annealed at 600 ° C. or higher 700 ° C. or less of the temperature range, {211} {100} intensity ratio of surface 0.7 or more with respect to surface and {211} with respect to plane {110} A method for manufacturing a rotor core of a motor, wherein the strength ratio of the surface is 0.4 or less . The method for manufacturing a rotor core of a motor according to claim 8, wherein the steel material has a component composition satisfying the following expression (1).
Record
0.50 ≦ (C / 12) / [(Ti / 48) + (Mo / 96)] ≦ 1.50 (1) The steel material is further
Nb: 0.08 mass% or less,
The method for producing a rotor core of a motor according to claim 8 or 9, comprising one or more selected from V: 0.15 mass% or less and W: 1.5 mass% or less. The method for manufacturing a rotor core of a motor according to claim 10, wherein the steel material has a component composition satisfying the following expression (2).
Record
0.50 ≦ (C / 12) / [(Ti / 48) + (Mo / 96) + (Nb / 93) + (V / 51) + (W / 184)] ≦ 1.50 (2) The steel material further includes S: 0.01 mass% or more and 0.1 mass% or less, and
Pb: 0.2 mass% or less,
Ca: 0.005 mass% or less,
The method for producing a rotor core of a motor according to any one of claims 8 to 11, comprising one or more of Bi: 0.1 mass% or less and B: 0.02 mass% or less.

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