JP4660474B2 - Non-oriented electrical steel sheet with excellent punching workability and magnetic properties after strain relief annealing and its manufacturing method - Google Patents

Description

  TECHNICAL FIELD The present invention relates to a non-oriented electrical steel sheet used as an iron core material for electrical equipment and a method for producing the non-oriented electrical steel sheet, and particularly to a non-oriented electrical steel sheet excellent in punching workability and magnetic properties after strain relief annealing and the production thereof. Regarding the method.

  In recent years, due to the increasing energy savings in electrical equipment worldwide, non-oriented electrical steel sheets used as the core material of rotating machines are required to have higher performance characteristics. Higher grade materials with increased Si and Al content to increase specific resistance and larger crystal grain size have come to be used. On the other hand, performance improvement is also required for general-purpose motors, but it is difficult to switch to high-grade materials like high-efficiency models due to severe cost constraints.

  As a steel sheet required for general-purpose models, Si is 1.5% or less, and it is a material that dramatically improves iron loss by promoting crystal grain growth during strain relief annealing performed after motor core punching. . Furthermore, recently, an increasing number of consumers use scrap generated during core punching as a raw material for casting, and it has become necessary to make the Al content of steel sheets less than 0.05% from the viewpoint of securing the castability of scrap.

In order to improve grain growth during strain relief annealing, it is important to reduce or detoxify the precipitates inevitably mixed in the steel, but the material design of non-oriented electrical steel sheets with Al: less than 0.05% As is seen in Patent Document 1, one of them is Al deoxidized steel (Sol. Al: about 0.02%), and about 0.002% of B is added to produce BN as a nitride. This is a method for suppressing precipitation of AlN, which is harmful to grain growth.
On the other hand, as seen in Patent Document 2, the SiO ₂ and MnO ratio of the oxide contained in the Si-deoxidized steel (Sol.Al ≦ 0.001%) is controlled, and stretchable inclusions harmful to crystal grain growth are added. As a countermeasure, Patent Document 3 discloses that the Mn / Si of the steel component is controlled to 0.2 to 1.0.

  As a method for improving the magnetism by further reducing or detoxifying sulfides after appropriately controlling the nitride or oxide generated by deoxidation in this way, for example, Patent Document 4 obtains a magnetic improvement effect. Therefore, S is defined as 0.005% or less. In Patent Document 5, when S: 0.01 to 0.02% is contained, Ca or rare earth element is added to fix S. In Patent Document 6, S is 0.0050%. A method for preventing fine precipitation during hot rolling of MnS by setting the heating temperature of the slab to 1100 ° C. or lower is disclosed below.

JP 54-163720 A Japanese Unexamined Patent Publication No. 7-150248 JP-A-6-73510 JP 58-117828 A JP 58-164724 A JP-A-4-63228

  In a situation where further reduction of iron loss is required, it has become difficult to manufacture sufficiently and stably by the above method. The present invention has been made in view of such problems, and B is added for the purpose of suppressing precipitation of AlN, and an improvement measure focusing on S, which is an unavoidable mixed element in steel, is obtained after strain relief annealing. An object of the present invention is to provide a steel sheet having good crystal grain growth, low iron loss, and high magnetic flux density.

The present invention has been made to solve the above problems, and the gist thereof is as follows.
(1) By mass%, Si: 1.5% or less, Mn: 0.4% or more and 1.5% or less, Sol. Al: 0.01% or more and 0.04% or less; Ti: 0.0015% or less; N: 0.0030% or less; S: 0.0010% or more and 0.0040% or less; 5 or more and 1.5 or less, consisting of the balance Fe and inevitable impurities, and 10% or more of the sulfides containing Mn are combined and precipitated with B precipitates, and MnS, Cu ₂ S and the composite sulfides are mixed. total distribution density is at 3.0 × 10 ⁵ cells / mm ² or less, wherein the distribution density of 1.0 × 10 ³ cells / mm ² or less der Rukoto of Ti precipitates of less than the diameter 0.1μm Non-oriented electrical steel sheet.
(2) and not more than 30μm grain size of the steel plate, non-oriented electrical steel sheet according to the grain size after stress relief annealing of 750 ° C. × 2 hours and wherein the at 50μm or more (1) .
( 3 ) Furthermore, Sn or Cu or Ni is used in a total of 0.01% or more and 0.50% or less of Sn, Cu or Ni, and / or one or more of REM, Ca or Mg. The non-oriented electrical steel sheet according to (1) or (2) , characterized by containing 0.001 to 0.1%.
( 4 ) As a method for producing the steel sheet according to any one of (1) to ( 3 ), in the case of finishing annealing following steelmaking, hot rolling, pickling, and cold rolling, hot slab heating is performed. A non-oriented electrical steel sheet characterized by being retained for 5 minutes or more in the range of 1150 ° C. or more and 1250 ° C. or less, and continuously hot rolling after 15 minutes or more in the range of 1050 ° C. or more and less than 1150 ° C. Production method.
( 5 ) The method for producing a non-oriented electrical steel sheet according to (4), wherein an exit temperature of the hot rolling finish rolling is set to 800 ° C. or higher.
(6) the hot-rolled finishing rolling exit temperature T (° C.), no set forth in, characterized in that the T ≧ 900-1000 × Sn [mass%] in steel sheet containing Sn (4) A method for producing grain-oriented electrical steel sheets.

According to the present invention, it is possible to provide a non-oriented electrical steel sheet having good crystal grain growth after strain relief annealing, low iron loss, and high magnetic flux density.

In the steel having Si of 1.5% or less, the present inventors reduced S, which is an unavoidable element, to about 0.0010%, and the slab heating temperature to 1100 ° C. or less as described in JP-A-4-632228. We faced the problem that iron loss after strain relief annealing does not stabilize even at low temperatures. When the cause was investigated, MnS, Cu ₂ S and its composite sulfide were finely dispersed in the steel in spite of the low S content and slab heating temperature. It was found that grain growth was remarkably suppressed. When observed in more detail, it was found that these sulfides were spherical with a diameter of about 0.1 to 0.3 μm, but contained Ti precipitates with a diameter equivalent to about 0.05 μm at the center. It has been found that the reason why the sulfide takes such a precipitation form is that TiN initially precipitated after casting to hot rolling slab heating is finely dispersed and the sulfide is precipitated in the core.

  In order to overcome this situation, the present inventors have focused on sulfides that have a higher growth rate than TiN and tend to coarsen. That is, contrary to the current situation where sulfides are precipitated together with TiN as a nucleus, it was found that as a result of attempting to precipitate TiN as a nucleus with sulfide as a nucleus, stable and low iron loss can be obtained. Further, the present invention has been completed based on the knowledge that a low iron loss can be stably obtained by compositely depositing sulfides for BN produced as the fixing of surplus N described later. Hereinafter, the experimental results that led to the present invention will be described.

(Experiment 1)
In a laboratory vacuum melting furnace, by mass%, C: 0.003%, Si: 0.6%, Mn: 0.1-0.8%, Sol.Al: 0.03%, Ti: 0.0012%, N: 0.0021%, S: 0.0005 Steel pieces containing ˜0.0025%, B: 0.0020%, Sn: 0.08% were prepared. These steel slabs are held at 1200 ° C for 20 minutes, cooled to 1100 ° C over 10 minutes, held for 30 minutes, then hot rolled to a sheet thickness of 2.5 mm, and then cold-rolled to 0.50 mm after pickling. did.
The cold-rolled sheet thus obtained was subjected to finish annealing at 800 ° C. for 10 seconds and then subjected to strain relief annealing at 750 ° C. for 2 hours, and the crystal grain size and iron loss were measured. As a result, as shown in Table 1, in samples 8, 9, 11, and 12 with Mn of 0.4% or more and S of 0.0012, 0.0025%, a good iron loss was obtained when the crystal grain size after strain relief annealing was 50 μm or more. It was.

Next, when the precipitates were observed for the samples after strain relief annealing, in the samples of 9, 9, 11, and 12, where good iron loss was obtained, the number ratio of sulfides containing Mn was 10% or more. It was observed that precipitates and composite precipitation occurred. On the other hand, in other samples having poor iron loss, a large number of fine precipitates less than 0.1 μm in diameter, which are considered to be B precipitates, were observed.
As described above, the crystal grain growth after the stress relief annealing and the iron loss improvement effect which are expressed by Mn: 0.4% or more and S: 0.0012, 0.0025% are considered as follows. First, the MnS precipitation start temperature rises by increasing Mn. As a result, MnS precipitates during the hot rolling prior to BN, and the BN that precipitates thereafter precipitates with MnS as a nucleus. Thereby, it is considered that the formation of fine precipitates of B can be suppressed, and crystal grain growth and iron loss after good strain relief annealing can be obtained.
On the other hand, if Mn is less than 0.4%, MnS is not sufficiently precipitated at the stage where BN precipitates. It is considered that the characteristics after strain relief annealing do not improve due to single and fine precipitation.

(Experiment 2)
In a laboratory vacuum melting furnace, by mass%, C: 0.0034%, Si: 0.75%, Mn: 0.15-0.72%, Sol.Al: 0.019%, Ti: 0.0008-0.0017%, N: 0.0018%, S A steel slab containing 0.0023%, B: 0.0025%, Sn: 0.03%, Cu: 0.01%, and Ni: 0.02% was produced. These steel slabs were held at 1200 ° C. for 5 minutes, lowered to 1100 ° C. and held for 30 minutes, then hot rolled to a sheet thickness of 2.7 mm, and cold-rolled to 0.50 mm after pickling. The cold-rolled sheet thus obtained was subjected to finish annealing at 800 ° C. for 10 seconds and then subjected to stress relief annealing at 750 ° C. for 2 hours, and the iron loss was measured while observing precipitates and crystal grain sizes of the steel sheet. .

As a result, as shown in Table 2, in samples 4, 5, 7, and 8 where Mn is 0.4% or more and Ti is 0.0015% or less, the sulfide density after strain relief annealing is 3.0 × 10 ⁵ or less, the average crystal The particle size was 50 μm or more, and good iron loss was obtained. In these samples, there were many spherical sulfides having a diameter of 0.2 to 0.3 μm, and a large number of fine Ti precipitates having a diameter of less than 0.1 μm on the outer periphery were confirmed. On the other hand, in Samples 1 to 3 with a low Mn of 0.15%, the sulfide density after strain relief annealing was high as 4.5 × 10 ⁵ pieces / mm ² or more, the average crystal grain size was as small as 35 μm or less, and the iron loss was bad. . The sulfides found in these samples were as small as 0.1 μm or less in diameter, and many sulfides were confirmed to contain Ti precipitates at the center. In addition, Mn is 0.4% or more, but for Samples 6 and 9 where Ti exceeds 0.0015%, the sulfide density after strain relief annealing is as high as 3.8 × 10 ⁵ pieces / mm ² or more, and the average grain size is also The iron loss was relatively bad at 45 μm or less. In these samples, sulfides varied in a wide range of diameters of 0.05 to 0.3 μm or less, and composite forms with Ti precipitates varied from the outer periphery or the center. Note that the distribution density of Ti precipitates corresponding to the diameter and less than 0.1 μm was as low as less than 1.0 × 10 ³ pieces / mm ^{2 in} all samples.

  Thus, the effect expressed by increasing Mn to 0.4% or more and Ti to 0.0015% or less is considered as follows. First, increasing Mn raises the MnS precipitation start temperature, while reducing Ti lowers the TiN precipitation start temperature. This reverses the order of precipitation from normal, and MnS is deposited before TiN. Next, MnS whose precipitation start temperature has increased precipitates and grows at 1200 ° C. before the hot rolling heating. On the other hand, it is considered that TiN having a reduced precipitation start temperature was precipitated at 1100 ° C. after the hot rolling and had already coarsened MnS.

(Experiment 3)
Next, in order to investigate the influence of heating conditions on hot rolling, using the steel pieces of Mn: 0.4% and S: 0.0025% in Experiment 1, hot rolling heating was held at 1200 ° C for 60 minutes and held at 1100 ° C for 60 minutes. Two levels were added, and after pickling, samples were prepared in the same process as in Experiment 1 for comparative evaluation. As a result, as shown in Table 3, Sample 1 (Sample 9 in Experiment 1), which had good results, had a composite precipitation ratio of almost zero in Sample 2 held at 1200 ° C, with a single and fine B precipitate. Were observed, and the crystal grain size and iron loss after strain relief annealing were the worst. On the other hand, in the sample 3 held at 1100 ° C., although not as high as the 1200 ° C. heating material (sample 2), the composite precipitation rate was as low as 5%, and neither the crystal grain size nor the iron loss was excellent. This result is considered as follows. First, when held at 1200 ° C, MnS is precipitated in combination with Mn being as high as 0.42%, but B is unprecipitated and is hot rolled, so it is single and fine during hot rolling and during strain relief annealing. B precipitates are formed, and grain growth and iron loss are remarkably deteriorated. Next, when it is held at 1100 ° C, the distribution of MnS becomes coarse and the number of nuclei from which BN precipitates is insufficient, so some B precipitates alone and finely, resulting in good crystal grain size and iron It is thought that the loss was not obtained.

(Experiment 4)
Next, in order to investigate the influence of the heating cycle of hot rolling, using the steel pieces of Mn: 0.42% and Ti: 0.0013% of Experiment 2, hot rolling heating was held at 1200 ° C for 60 minutes and held at 1100 ° C for 60 minutes. Two levels were added and hot-rolled, and after pickling, samples were prepared in the same process as in Experiment 2 for comparative evaluation. As a result, as shown in Table 4, sample 1 (sample 5 in experiment 2), which had good results, sample 2 held at 1200 ° C. had a high distribution density of Ti precipitates although the sulfide density was low, The crystal grain size and iron loss after strain relief annealing were the worst. On the other hand, Sample 3 maintained at 1100 ° C. had a low Ti precipitate density but a high sulfide density, and neither crystal grain size nor iron loss was excellent. This result is considered as follows. First, when held at 1200 ° C, the MnS coarsening progresses due to the high Mn of 0.42%. Fine Ti precipitates are produced, and grain growth and iron loss are significantly worsened. Next, when held at 1100 ° C., the growth of MnS is insufficient, and it is considered that the number of sulfides increased and good crystal grain size and iron loss were not obtained.

To summarize the above, the present invention preferentially and optimally precipitates MnS by optimizing the heating cycle of the amount of Mn and hot rolling, and simultaneously detoxifies both MnS and BN by compound precipitation of fine BN. The present inventors have found out measures for improving crystal grain growth and iron loss. In order to realize this, the amount of Al must be reduced while increasing the amount of Mn. This is because Al consumes N as AlN, and AlN itself hinders crystal grain growth. On the other hand, in the production method in which no Al is added, the amount of Al can be kept very low, which is convenient for the generation of BN. However, the inclusion of many SiO ₂ and MnO remaining in the steel is accompanied by an increase in the amount of Mn. Stretching, on the contrary, deteriorates grain growth. Accordingly, in the present invention, 0.01% to 0.04% was found as the optimum range of Al in which stretching is suppressed by modification of inclusions and BN precipitation is prioritized.

In addition, the present invention optimizes the deposition temperature of MnS and TiN, optimizes the hot rolling heating cycle, first precipitates MnS coarsely, and then compositely precipitates fine TiN. They found ways to detoxify and improve grain growth and iron loss.
In order to realize this, it is necessary to increase the Mn content to increase the MnS precipitation temperature, and further reduce the Al content to suppress the formation of AlN and promote the TiN precipitation. As a method for keeping the amount of Al extremely low, there is a method of using Si for deoxidation in steelmaking. In this case, it was found that inclusions are elongated when the Mn amount is increased, and the iron loss is worsened. Therefore, in the present invention, by controlling the Al amount within a relatively small range of 0.01 to 0.04%, it is possible to increase the Mn amount by modifying the inclusions while maintaining the precipitation of TiN. Further, in this Al content range, excess N after TiN precipitation tends to cause fine precipitation of AlN that is harmful to grain growth, so this was avoided by adding a small amount of B to form BN.

  Such a technical idea was first discovered in the present invention.For example, in Japanese Patent Laid-Open No. 58-117828, Si: 0.1 to 1.0%, Al: less than 0.1%, Mn: 0.75 to 1.5%, N / B: Although the content of B in the range of 0.7 to 1.2 is specified, there is no specification for the composite precipitation of MnS and BN, and the amount of Sol.Al to realize it and the heating temperature of hot rolling. The technical ideas are completely different, and the present invention cannot be analogized. In addition, in Japanese Patent Laid-Open No. 2000-248344, Si: 1.8% or less, Sol.Al: 0.05-0.20%, Mn: 0.05-1.5% are specified, but if Sol.Al: 0.04% is exceeded, rather than BN Since AlN is preferentially precipitated, the technical idea of the present invention of precipitating BN with MnS as a nucleus does not hold.

Next, the reasons for limiting the numerical values of the components and products in the present invention will be described.
Si is an effective element for increasing the electrical resistance, but if added over 1.5%, the hardness is increased, the magnetic flux density is decreased, and the cost is increased.
Mn is an important element for expressing the present invention. The main purpose of the present invention is to deposit BN and / or TiN with a sulfide containing MnS as a nucleus. For this purpose, MnS must be sufficiently precipitated before the precipitation temperature of BN and / or TiN. In the present invention containing Ti: 0.0015% or less, N: 0.0030% or less, and B / N of 0.5 or more and 1.5 or less, this object is achieved by making Mn 0.4% or more. If the addition exceeds 1.5%, the saturation magnetic flux density is significantly lowered, and the γ → α transformation temperature is lowered to make it difficult to control the structure of the hot-rolled sheet, so 1.5% was made the upper limit.

Al is an element necessary for deoxidation of steel. If less than 0.01% Sol, undeoxidized oxygen remains in the steel to form SiO ₂ · MnO oxide, which is combined with the effect of Mn added at 0.4% or more to stretch and grow grains Therefore, the lower limit of Sol.Al was set to 0.01%. In addition, if Sol.Al exceeds 0.04%, AlN will be deposited instead of BN, and the expression of the present invention will be difficult.In addition, from the viewpoint of securing TiN precipitation and scrap utilization at the consumer, The upper limit of Sol.Al was 0.04%.

Ti produces TiN and remarkably deteriorates the grain growth, but it is an unavoidable element, so it is industrially difficult to make it zero. In the present invention, the upper limit is defined as 0.0015% as an allowable amount that can be rendered harmless by composite precipitation with MnS, Cu ₂ S, and composite sulfides thereof. If Ti exceeds 0.0015%, the TiN precipitation start temperature rises and the preferential precipitation of MnS cannot be controlled.

  N generates TiN and AlN in addition to BN. In the present invention containing Sol. Al: 0.01 to 0.04%, when AlN is generated, crystal grain growth is remarkably deteriorated. Therefore, it is necessary to add B to suppress the generation of AlN. Therefore, the amount of B to be added must be increased as N increases. However, excessive addition of B causes embrittlement of the steel sheet and deteriorates productivity, so the upper limit of N was set to 0.0030%.

S is necessary for producing a sulfide which becomes a precipitation nucleus of BN and / or TiN, and the object of the present invention can be achieved by containing 0.0010% or more. However, if it exceeds 0.0040%, the amount of sulfide precipitates itself increases and the grain growth is inhibited, so 0.0040% was made the upper limit.
B is an element that must be added in order to suppress the formation of AlN that is harmful to crystal grain growth, but for that purpose, it is necessary to add 0.5 or more as B / N. Even if it is added excessively with respect to N, the effect is saturated, so the upper limit was made 1.5 for B / N.

Sn, Cu, and Ni are effective in suppressing nitriding and oxidation of the steel sheet surface during annealing, especially strain relief annealing, and in the steel of the present invention containing Sol. It is preferable to add. If the amount added is less than 0.01%, there is no effect, and even if added over 0.50%, the effect is saturated and the cost is increased, so the range of the amount added is set to 0.01% to 0.50%. In addition, since the nitriding / oxidation suppressing effects of Sn, Cu, and Ni are the same, the above addition range may be satisfied by a single or a composite. In addition, it is also possible to add 0.001 to 0.5% of one or more of REM, Ca, and Mg.
In particular, Sn is an extremely effective element for improving the magnetic flux density in the present invention. This is because in the present invention, since the Mn is increased, the γ → α transformation temperature is inevitably lowered, and the grain growth of the hot-rolled sheet cannot be sufficiently promoted, and this needs to be compensated. If the amount added is less than 0.01%, there is no effect. If the amount added exceeds 0.50%, the effect is saturated and the cost increases. Therefore, the range of the amount added is set to 0.01% to 0.50%. Furthermore, Sn also has an effect of suppressing nitriding and oxidation of the steel sheet surface during strain relief annealing, and addition from the viewpoint is also desirable.

For the composite precipitation characterized by the present invention, the number ratio of the sulfide containing Mn in which the B precipitate is compositely precipitated is defined as 10% or more. This is based on the observation result in a sample having no single and fine B precipitate.
The crystal grain size is an important factor for achieving both punching workability and magnetism. In the steel sheet to be subjected to punching, the grain size is set to 30 μm or less because the punching processability deteriorates when the particle size exceeds 30 μm. Moreover, since the required iron loss is not satisfied unless the grain size is less than 50 μm in electrical products, the crystal grain size after 750 ° C. × 2 hours of strain relief annealing, which is generally performed, is defined as 50 μm or more.

If there is too much MnS, Cu ₂ S and its composite sulfide, the grain growth is inhibited. In order to obtain a crystal grain size of 50 μm or more after strain relief annealing, the distribution density must be 3.0 × 10 ⁵ pieces / mm ² or less. The distribution density described here refers to the number of precipitates in which Mn, S or Cu, S, or Mn, Cu, S is detected by observing a chemically polished sample after mirror polishing with a scanning or transmission electron microscope. It is divided by the observation visual field area (the total area when a plurality of visual fields are observed).

Ti precipitates can be rendered harmless even though they are fine enough to be less than 0.1 μm in diameter by compound precipitation of sulfides in the nucleus. The crystal grain size of 50 μm or more was obtained after strain relief annealing, and the iron loss was good when the distribution density of Ti precipitates with diameter equivalent to less than 0.1 μm was 1.0 × 10 ³ pieces / mm ² or less Therefore, this was made the upper limit.

Next, the reasons for limiting the manufacturing conditions in the present invention will be described.
Hot rolling slab heating needs to be a two-stage continuous cycle because BN and / or TiN is deposited and harmed by using sulfides containing MnS, Cu ₂ S and their composite sulfides as nuclei. In the steel of the present invention in which Mn is increased to 0.4% or more, precipitation and growth of MnS become prominent in the temperature range of 1150 ° C or higher, but solid solution proceeds when the temperature exceeds 1250 ° C. The temperature was 1150 ° C or higher and 1250 ° C or lower. Since the growth rate of MnS is fast, it is sufficient that the residence time in this temperature range is 5 minutes or more. Next, since TiN and BN proceeded with compound precipitation on MnS at a temperature of less than 1150 ° C., the subsequent heating temperature was set to less than 1150 ° C., and the lower limit temperature was set to 1050 ° C. from the viewpoint of securing rollability. In addition, although the precipitation of TiN proceeds when the subsequent heating temperature is lower than 1050 ° C., the compounding with sulfides is insufficient, so that single and fine precipitates increase during hot rolling. The latter heating time was set to 15 minutes or more in consideration of the precipitation time of TiN and BN. More preferably, it is 30 minutes or more.

  The exit temperature of hot rolling finish rolling is higher at 800 ° C or higher as much as possible at a temperature of γ → α transformation or less, but the magnetic flux density becomes higher, but the temperature is relaxed by the amount of Sn added, so Sn When T was added, it was specified that T ≧ 900-1000 × Sn [mass%] in consideration of the degree of relaxation.

  In a laboratory vacuum melting furnace, by mass%, C: 0.003%, Si: 0.55%, Mn: 0.12-0.96%, Sol.Al: 0.033%, Ti: 0.0008%, N: 0.0025%, S: 0.0032 Steel pieces containing%, B: 0.0017%, Sn: 0.02 to 0.09% were produced. These steel pieces were heated to 1230 ° C. and held for 10 minutes, then cooled to 1120 ° C. and held for 30 minutes, and then hot rolled to a thickness of 2.5 mm. The exit temperature of finish rolling was 855 ° C. The hot-rolled sheet was pickled and cold-rolled to a thickness of 0.50 mm, then subjected to finish annealing at 825 ° C. for 10 seconds, and then subjected to strain relief annealing at 750 ° C. for 2 hours. The thus obtained sample was measured for crystal grain size, iron loss, and magnetic flux density, and the precipitate was observed with a transmission electron microscope. As a result, as shown in Table 5, in samples 7 to 12 where Mn is 0.4% or more, good iron loss is obtained when the average crystal grain size after strain relief annealing is 50 μm or more, and the number ratio of composite precipitation is also 10%. That was all. Furthermore, samples 8, 9, 11, and 12 satisfying 855 ≧ 900-1000 × Sn gave a high magnetic flux density of about 0.02T.

  In a laboratory vacuum melting furnace, in mass%, C: 0.003%, Si: 1.3%, Mn: 0.29 to 1.08%, Al: 0.027%, Ti: 0.0013%, N: 0.0019%, S: 0.0026%, Steel pieces containing B: 0.0024% and Sn: 0.07% were produced. These steel slabs were heated to 1230 ° C, immediately cooled to 1090 ° C, held for 30 minutes, and then hot rolled. In this heating, the steel slab stayed at a temperature of 1150 ° C. or higher for 15 minutes, and the finish rolling exit temperature was 840 ° C. In addition, a test was also conducted in which hot rolling was performed immediately after heating at a constant temperature of 1230 ° C. or 1090 ° C. for 60 minutes. The hot-rolled sheet thus obtained was pickled, cold-rolled to a thickness of 0.50 mm, subjected to finish annealing at 850 ° C. for 10 seconds, subjected to stress relief annealing at 750 ° C. for 2 hours, The iron loss and magnetic flux density were measured, and the precipitate was observed with a transmission electron microscope. As a result, as shown in Table 6, in Samples 10 to 12 in which the Mn is 0.4% or more and the hot rolling temperature is 1230 → 1090 ° C. in a two-stage cycle, the average grain size after strain relief annealing is 50 μm or more. Good iron loss was obtained, and the number ratio of composite precipitates was 10% or more. Furthermore, all satisfied 840 ≧ 900-1000 × Sn (= 0.07%), and a high magnetic flux density was obtained.

  In a laboratory vacuum melting furnace, by mass%, C: 0.0038%, Si: 0.51%, Mn: 0.12-0.84%, Sol.Al: 0.025%, Ti: 0.0008-0.0024%, N: 0.0025%, S A steel slab containing 0.0035% and 0.0016% B was produced. These steel slabs were heated to 1240 ° C, immediately cooled to 1120 ° C, held for 30 minutes, and then hot rolled to a thickness of 2.7 mm. In this heating, the steel slab stayed at a temperature of 1150 ° C or higher was 22 minutes, and the finish rolling exit temperature was 820 ° C. This hot-rolled sheet is pickled and cold-rolled to a thickness of 0.50 mm, then subjected to finish annealing at 825 ° C for 10 seconds, and then subjected to strain-relief annealing at 750 ° C for 2 hours. The iron loss was measured while observing.

As a result, as shown in Table 7, in Samples 7, 8, 10, and 11 where Mn is 0.4% or more and Ti is 0.0015% or less, the sulfide density after strain relief annealing is 3.0 × 10 ⁵ or less, the average crystal Good iron loss was obtained when the particle size was 50 μm or more. In these samples, there were many spherical sulfides having a diameter of 0.2 to 0.3 μm, and a large number of Ti precipitates were observed on the outer periphery of the sulfide. On the other hand, in Samples 1 to 6 having a low Mn of 0.12 and 0.25%, the sulfide density after strain relief annealing was high, the average crystal grain size was small, and the iron loss was poor. The sulfides found in these samples were as small as 0.1 μm or less in diameter, and it was confirmed that a large number of sulfides contained fine Ti precipitates. In addition, Mn was 0.4% or more, but Samples 9 and 12 in which Ti exceeded 0.0015% also had a high sulfide density after strain relief annealing, a small average crystal grain size, and a relatively low iron loss. In these samples, sulfides varied in a wide range of 0.05 to 0.3 μm or less in diameter, and composite forms with Ti precipitates varied from the periphery or the center of the sulfides. The distribution density of Ti precipitates with a diameter of less than 0.1 μm was as low as less than 1.0 × 10 ³ pieces / mm ^{2 in} all samples.

In a laboratory vacuum melting furnace, by mass%, C: 0.0022%, Si: 1.2%, Mn: 0.31-1.44%, Sol.Al: 0.03%, Ti: 0.0013%, N: 0.0016%, S: 0.0031 Steel pieces containing%, B: 0.0021%, Sn: 0.02% were produced. These steel slabs were heated to 1220 ° C, immediately cooled to 1070 ° C, held for 20 minutes, and then hot rolled. In this heating, the time during which the steel slab stayed at a temperature of 1150 ° C. or higher was 15 minutes. In addition to this, a test was also conducted in which hot rolling was performed immediately after heating at a constant temperature of 1220 ° C. or 1070 ° C. for 45 minutes. The hot-rolled sheet thus obtained is pickled, cold-rolled to a thickness of 0.50 mm, then subjected to finish annealing at 850 ° C. for 5 seconds, and then subjected to strain relief annealing at 750 ° C. for 2 hours. The iron loss was measured while observing the particle size. As a result, as shown in Table 8, in samples 12 to 15 where Mn is 0.4% or more and the heating temperature of hot rolling is 1220 → 1070 ° C., the sulfide density after strain relief annealing is 3.0 × 10 ⁵ or less, the average The crystal grain size was 50 μm or more, and good iron loss was obtained. In these samples, there were many spherical sulfides having a diameter of 0.2 to 0.3 μm, and a large number of Ti precipitates were observed on the outer periphery of the sulfide. In other samples, the sulfide density after strain relief annealing was high, or the Ti precipitate density was high, the average crystal grain size was small, and the iron loss was bad.

Claims (6)

In terms of mass%, Si: 1.5% or less, Mn: 0.4% or more and 1.5% or less, Sol. Al: 0.01% or more and 0.04% or less; Ti: 0.0015% or less; N: 0.0030% or less; S: 0.0010% or more and 0.0040% or less; 5 or more and 1.5 or less, consisting of the balance Fe and inevitable impurities, and 10% or more of the sulfides containing Mn are combined and precipitated with B precipitates, and MnS, Cu ₂ S and the composite sulfides are mixed. total distribution density is at 3.0 × 10 ⁵ cells / mm ² or less, wherein the distribution density of 1.0 × 10 ³ cells / mm ² or less der Rukoto of Ti precipitates of less than the diameter 0.1μm Non-oriented electrical steel sheet. The non-oriented electrical steel sheet according to claim 1, wherein the crystal grain size of the steel sheet is 30 μm or less and the crystal grain size after strain relief annealing at 750 ° C. × 2 hours is 50 μm or more. Further, one or more of Sn, Cu, and Ni are added in a total mass of 0.01% to 0.50%, and / or one or more of REM, Ca, and Mg are 0.001. The non-oriented electrical steel sheet according to claim 1 or 2 , characterized by comprising ~ 0.1%. As a method for producing the steel sheet according to any one of claims 1 to 3 , the hot-rolling slab heating is performed at 1150 ° C. or more and 1250 at the time of finishing annealing subsequent to steelmaking, hot rolling, pickling and cold rolling. A method for producing a non-oriented electrical steel sheet, wherein the steel sheet is retained for 5 minutes or more in a range of ℃ or less, and continuously hot rolled for 15 minutes or more in a range of 1050 ° C. or more and less than 1150 ° C. The method for producing a non-oriented electrical steel sheet according to claim 4, wherein an outlet temperature of the hot rolling finish rolling is set to 800 ° C. or more. The non-directional electromagnetic according to claim 4 , wherein an exit temperature T (° C.) of the hot rolling finish rolling is T ≧ 900−1000 × Sn [mass%] in a steel plate containing Sn. A method of manufacturing a steel sheet.