JP4272557B2 - Method for producing unidirectional electrical steel sheet with excellent magnetic properties - Google Patents
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- 238000004519 manufacturing process Methods 0.000 title claims description 21
- 229910000976 Electrical steel Inorganic materials 0.000 title claims description 16
- 238000000137 annealing Methods 0.000 claims description 80
- 238000001953 recrystallisation Methods 0.000 claims description 55
- IJGRMHOSHXDMSA-UHFFFAOYSA-N Atomic nitrogen Chemical compound N#N IJGRMHOSHXDMSA-UHFFFAOYSA-N 0.000 claims description 36
- 238000005121 nitriding Methods 0.000 claims description 33
- 238000005097 cold rolling Methods 0.000 claims description 23
- 238000005098 hot rolling Methods 0.000 claims description 22
- 238000005261 decarburization Methods 0.000 claims description 21
- 229910052757 nitrogen Inorganic materials 0.000 claims description 21
- 229910000831 Steel Inorganic materials 0.000 claims description 18
- 239000010959 steel Substances 0.000 claims description 18
- QGZKDVFQNNGYKY-UHFFFAOYSA-N Ammonia Chemical compound N QGZKDVFQNNGYKY-UHFFFAOYSA-N 0.000 claims description 10
- 229910052739 hydrogen Inorganic materials 0.000 claims description 8
- 239000001257 hydrogen Substances 0.000 claims description 8
- 238000001556 precipitation Methods 0.000 claims description 8
- 238000001816 cooling Methods 0.000 claims description 6
- 229910021529 ammonia Inorganic materials 0.000 claims description 5
- 239000007789 gas Substances 0.000 claims description 5
- 238000005096 rolling process Methods 0.000 claims description 5
- 229910052718 tin Inorganic materials 0.000 claims description 5
- 229910052787 antimony Inorganic materials 0.000 claims description 4
- 229910001224 Grain-oriented electrical steel Inorganic materials 0.000 claims description 2
- 239000012535 impurity Substances 0.000 claims description 2
- 230000009467 reduction Effects 0.000 claims description 2
- 125000004435 hydrogen atom Chemical class [H]* 0.000 claims 1
- 239000003112 inhibitor Substances 0.000 description 53
- 238000000034 method Methods 0.000 description 49
- 238000010438 heat treatment Methods 0.000 description 40
- 230000000694 effects Effects 0.000 description 15
- 239000000126 substance Substances 0.000 description 11
- 239000006104 solid solution Substances 0.000 description 10
- 239000012298 atmosphere Substances 0.000 description 8
- 239000011521 glass Substances 0.000 description 8
- 230000015572 biosynthetic process Effects 0.000 description 7
- XEEYBQQBJWHFJM-UHFFFAOYSA-N Iron Chemical group [Fe] XEEYBQQBJWHFJM-UHFFFAOYSA-N 0.000 description 6
- 238000009776 industrial production Methods 0.000 description 6
- 239000002245 particle Substances 0.000 description 6
- 230000008569 process Effects 0.000 description 6
- UFHFLCQGNIYNRP-UHFFFAOYSA-N Hydrogen Chemical compound [H][H] UFHFLCQGNIYNRP-UHFFFAOYSA-N 0.000 description 4
- 229910052782 aluminium Inorganic materials 0.000 description 4
- 239000010949 copper Substances 0.000 description 4
- 230000007423 decrease Effects 0.000 description 4
- 230000004907 flux Effects 0.000 description 4
- 239000002244 precipitate Substances 0.000 description 4
- 239000006185 dispersion Substances 0.000 description 3
- 150000002431 hydrogen Chemical class 0.000 description 3
- 238000002791 soaking Methods 0.000 description 3
- 229910052717 sulfur Inorganic materials 0.000 description 3
- ILRRQNADMUWWFW-UHFFFAOYSA-K aluminium phosphate Chemical compound O1[Al]2OP1(=O)O2 ILRRQNADMUWWFW-UHFFFAOYSA-K 0.000 description 2
- 230000008859 change Effects 0.000 description 2
- 239000003795 chemical substances by application Substances 0.000 description 2
- 239000011248 coating agent Substances 0.000 description 2
- 238000000576 coating method Methods 0.000 description 2
- 229910052802 copper Inorganic materials 0.000 description 2
- 230000006698 induction Effects 0.000 description 2
- 229910052742 iron Inorganic materials 0.000 description 2
- 239000010410 layer Substances 0.000 description 2
- 239000000463 material Substances 0.000 description 2
- 238000002844 melting Methods 0.000 description 2
- 230000008018 melting Effects 0.000 description 2
- 239000012299 nitrogen atmosphere Substances 0.000 description 2
- 239000011164 primary particle Substances 0.000 description 2
- 238000000746 purification Methods 0.000 description 2
- 238000005204 segregation Methods 0.000 description 2
- 229910052711 selenium Inorganic materials 0.000 description 2
- 238000000926 separation method Methods 0.000 description 2
- 239000000243 solution Substances 0.000 description 2
- XLYOFNOQVPJJNP-UHFFFAOYSA-N water Substances O XLYOFNOQVPJJNP-UHFFFAOYSA-N 0.000 description 2
- RYGMFSIKBFXOCR-UHFFFAOYSA-N Copper Chemical compound [Cu] RYGMFSIKBFXOCR-UHFFFAOYSA-N 0.000 description 1
- ATJFFYVFTNAWJD-UHFFFAOYSA-N Tin Chemical compound [Sn] ATJFFYVFTNAWJD-UHFFFAOYSA-N 0.000 description 1
- 238000009825 accumulation Methods 0.000 description 1
- 230000002411 adverse Effects 0.000 description 1
- 229910052804 chromium Inorganic materials 0.000 description 1
- 238000009749 continuous casting Methods 0.000 description 1
- 230000007547 defect Effects 0.000 description 1
- 238000000265 homogenisation Methods 0.000 description 1
- 238000012423 maintenance Methods 0.000 description 1
- 229910052748 manganese Inorganic materials 0.000 description 1
- 239000000155 melt Substances 0.000 description 1
- 238000002156 mixing Methods 0.000 description 1
- 229910052759 nickel Inorganic materials 0.000 description 1
- 150000004767 nitrides Chemical class 0.000 description 1
- 230000003647 oxidation Effects 0.000 description 1
- 238000007254 oxidation reaction Methods 0.000 description 1
- 229910052698 phosphorus Inorganic materials 0.000 description 1
- 229920006395 saturated elastomer Polymers 0.000 description 1
- 239000013589 supplement Substances 0.000 description 1
- 239000002344 surface layer Substances 0.000 description 1
- 229910052719 titanium Inorganic materials 0.000 description 1
Description
本発明は、主にトランス等の鉄芯として使用される一方向性電磁鋼板を極めて安定的に製造する方法に関するものである。 The present invention relates to a method for producing a unidirectional electrical steel sheet mainly used as an iron core of a transformer or the like very stably.
磁束密度Ｂ8（８００Ａ／ｍの磁場中での磁束密度）が１．９Ｔを越える、磁気特性の優れた一方向製電磁鋼板を安定的に生産を行う技術は種々提案されているが、Ａｌインヒビターとして含有する場合の製造方法はスラブ加熱温度により、表１に示す三種類の技術に分類できる。 Various techniques for stably producing a unidirectional electrical steel sheet having excellent magnetic properties with a magnetic flux density B 8 (magnetic flux density in a magnetic field of 800 A / m) exceeding 1.9 T have been proposed. The manufacturing method in the case of containing as an inhibitor can be classified into three types of techniques shown in Table 1 depending on the slab heating temperature.
第一の技術は、完全固溶非窒化型と呼ばれ、スラブを１３５０℃から最高では１４５０℃の超高温度に加熱し、かつ、スラブ全体を通して一様に加熱（均熱）するために十分な時間スラブをその温度に保持する方法である。これはＭｎＳ、ＡｌＮ等のインヒビター能力を有する物質を完全溶体化させて、二次再結晶に必要なインヒビターとして機能させるためのものであり、この完全溶体化処理は同時に、スラブ部位によるインヒビター強度差を解消する手段にもなっており、この点では、安定した二次再結晶発現に有利である。 The first technique, called the fully solid solution non-nitrided type, is sufficient to heat the slab from 1350 ° C to an extremely high temperature of up to 1450 ° C and uniformly (uniform temperature) throughout the slab. This is a method of keeping the slab at that temperature for a long time. This is to completely dissolve a substance having inhibitor ability such as MnS and AlN to function as an inhibitor necessary for secondary recrystallization, and this complete solution treatment is performed simultaneously with the difference in inhibitor strength depending on the slab site. In this respect, it is advantageous for stable secondary recrystallization.
しかしながらこの技術の場合、二次再結晶に必要なインヒビター量を確保するための完全溶体化温度は熱力学的にはあまり高くはないが、実際の工業生産では生産性とスラブ全体の均一固溶状態とを確保するため超高温度とならざるを得ず、これに伴い実生産において様々な問題を包含している。例えば、１）所定熱延温度の確保が困難であり、確保出来なかった場合、インヒビター強度のスラブ内偏差が生じるため二次再結晶不良が発生する、２）熱延加熱時に粗大粒が生成し、その粗大粒部分は二次再結晶出来ずに、線状の二次再結晶不良が発生する、３）スラブ表層が溶融しノロとなり加熱炉のメンテナンスに多大の労力が必要である、４）熱延後の鋼帯に巨大なエッジクラックが発生し歩留まりが低下する、等である。 However, in this technique, the complete solution temperature for securing the amount of inhibitor necessary for secondary recrystallization is not so high in thermodynamics, but in actual industrial production, productivity and uniform solid solution of the entire slab are obtained. In order to ensure the state, the temperature has to be extremely high, and accordingly, various problems are included in actual production. For example, 1) It is difficult to ensure a predetermined hot rolling temperature, and if it cannot be ensured, a secondary recrystallization failure occurs due to a deviation in the slab of inhibitor strength. 2) Coarse grains are formed during hot rolling heating. The coarse grains cannot be recrystallized, and linear secondary recrystallization failure occurs. 3) The slab surface melts and becomes noro, and a great deal of labor is required for the maintenance of the heating furnace. 4) A huge edge crack is generated in the steel strip after hot rolling, and the yield is lowered.
また、この技術では非特許文献１に開示されているように、インヒビターを補うため脱炭焼鈍後二次再結晶開始までに窒化処理を行うと、かえって磁気特性が劣化してしまう。 Further, as disclosed in Non-Patent Document 1, in this technique, if nitriding is performed before decarburization annealing and the start of secondary recrystallization in order to supplement the inhibitor, the magnetic characteristics are deteriorated.
第二の技術は、（完全）析出窒化型と呼ばれるもので、特許文献１〜３などに開示されているように、スラブ加熱温度を１２８０℃未満で行い、脱炭焼鈍後二次再結晶開始までに窒化処理を行うものである。このような方法においては、例えば特許文献４に示されるように脱炭焼鈍後の一次再結晶粒の平均粒径を一定範囲、通常１８〜３５μｍの範囲に制御することが、二次再結晶を良好に行わせる上で非常に重要である。 The second technique is called a (complete) precipitation nitriding type. As disclosed in Patent Documents 1 to 3, the slab heating temperature is less than 1280 ° C., and secondary recrystallization is started after decarburization annealing. The nitriding treatment is performed up to this point. In such a method, for example, as shown in Patent Document 4, it is possible to control the average grain size of primary recrystallized grains after decarburization annealing within a certain range, usually in the range of 18 to 35 μm. It is very important for good performance.
また、この技術ではスラブ加熱時の固溶窒素などインヒビター能力を有する物質の鋼中固溶量が一次再結晶粒成長性に大きく影響することから、スラブ内一次再結晶粒の大きさを均一にするため、例えば特許文献５では、スラブ加熱時の固溶窒素を低くして、後工程で生じる不均一な析出を抑制する方法が開示されている。そして固溶量低減の面から、実際のスラブ加熱温度は１１５０℃以下が望まれている。 In addition, in this technology, the amount of solid solution in the steel of a substance having inhibitor ability such as solute nitrogen during slab heating greatly affects the primary recrystallized grain growth, so the size of primary recrystallized grains in the slab is made uniform. For this reason, for example, Patent Document 5 discloses a method in which solid solution nitrogen during slab heating is lowered to suppress non-uniform precipitation that occurs in a subsequent process. From the viewpoint of reducing the amount of solid solution, the actual slab heating temperature is desired to be 1150 ° C. or lower.
しかしながら、この技術ではいかに厳密に成分を調整してもインヒビター物質を完全に析出させることは出来ないことから、実際の生産活動では一次再結晶粒径を一定にするため一次再結晶焼鈍の条件（特に温度）をコイル毎に調節している。このため製造工程は煩雑化し、また脱炭焼鈍後の酸化層形成が一定でないため、グラス皮膜形成の不安定化を生じる場合がある。 However, with this technique, no matter how strictly the components are adjusted, the inhibitor substance cannot be precipitated completely. Therefore, in actual production activities, the primary recrystallization annealing conditions ( Especially the temperature) is adjusted for each coil. For this reason, the manufacturing process becomes complicated, and since the formation of the oxide layer after decarburization annealing is not constant, the glass film formation may become unstable.
第三の技術は混合型と呼ばれ、実操業的には特許文献６に示すように、スラブ加熱温度を１２００〜１３５０℃とし、第二の技術と同様に窒化を必須とするものである。第一の技術における１３５０℃を超える超高温度のスラブ加熱温度を避けるため、インヒビター物質の初期含有量を減じてスラブ加熱温度を下げ、これに伴い不足するインヒビター強度を窒化処理により補充する。この技術はさらに２種類に分類される。一つは部分固溶窒化型（部分析出窒化型）、もう一つは特許文献７に代表される完全固溶窒化型である。 The third technique is called a mixed type. In actual operation, as shown in Patent Document 6, the slab heating temperature is set to 1200 to 1350 ° C., and nitriding is essential as in the second technique. In order to avoid the slab heating temperature exceeding 1350 ° C. in the first technique, the slab heating temperature is lowered by reducing the initial content of the inhibitor substance, and the insufficient inhibitor strength is supplemented by nitriding treatment. This technology is further classified into two types. One is a partial solid solution nitriding type (partial precipitation nitriding type), and the other is a complete solid solution nitriding type represented by Patent Document 7.
この技術では、インヒビターを一次インヒビターと二次インヒビターと区別している。ここで、一次インヒビターは主に一次再結晶粒径を決定するものであるが、もちろんこれも二次再結晶にも寄与している。二次インヒビターは二次再結晶を可能ならしめるものである。一次インヒビターの存在により、一次再結晶後の粒径変動が小さくなるため、一次再結晶焼鈍条件を粒径のために変更する必要がなく、グラス皮膜形成が極めて安定している。 This technique distinguishes inhibitors from primary inhibitors and secondary inhibitors. Here, the primary inhibitor mainly determines the primary recrystallization particle diameter, but of course this also contributes to the secondary recrystallization. Secondary inhibitors make secondary recrystallization possible. The presence of the primary inhibitor reduces the particle size variation after the primary recrystallization, so there is no need to change the primary recrystallization annealing condition for the particle size, and the glass film formation is extremely stable.
一次インヒビターとしては、第一の技術で用いられているインヒビター物質（例えばＡｌＮ，ＭｎＳ，ＭｎＳｅ，Ｃｕ−Ｓ，Ｓｎ，Ｓｂ）が用いられる。ただし、スラブ加熱温度を低減するためその含有量は少ないことが求められる。二次インヒビターは一次インヒビターと脱炭焼鈍後二次再結晶開始までで窒化され形成されたＡｌＮである。また特許文献７には一次インヒビターとしてその他にＢＮが記載されているが、ＮはＡｌとも結合するので実際的にはＡｌとＢを同時に含有すると二次再結晶が不安定になる場合がある。 As the primary inhibitor, an inhibitor substance (for example, AlN, MnS, MnSe, Cu—S, Sn, Sb) used in the first technique is used. However, in order to reduce the slab heating temperature, the content is required to be small. The secondary inhibitor is AlN formed by nitriding from the primary inhibitor to the start of secondary recrystallization after decarburization annealing. In Patent Document 7, BN is also described as a primary inhibitor. However, since N also binds to Al, when it contains Al and B at the same time, secondary recrystallization may become unstable.
前記三つの技術に共通の課題として、必要なインヒビター物質（特にＡｌとＮ）の含有量の適正範囲が狭いことがあげられる。そこで従来より、酸可溶性Ａｌ（以下ｓＡｌ）からＮ当量を控除したＡｌＲを指標として製造条件を調節する方法が、第一と第二の技術において開示されている。 A problem common to the above three techniques is that the appropriate range of the content of necessary inhibitor substances (especially Al and N) is narrow. Therefore, conventionally, methods for adjusting production conditions using AlR obtained by subtracting N equivalents from acid-soluble Al (hereinafter referred to as sAl) as an index have been disclosed in the first and second techniques.
第一の技術では、例えば特許文献８には、ＡｌＲ値によって、最終冷延前焼鈍の均熱時間もしくは冷却速度の他、一連の工程条件のうちいずれかを調節することを規定している。 In the first technique, for example, Patent Document 8 stipulates that any one of a series of process conditions is adjusted in addition to the soaking time or cooling rate of the annealing before the final cold rolling depending on the AlR value.
また第二の技術では、特許文献９には仕上焼鈍時の雰囲気中のＮ2の割合をＡｌＲの式により規定している。特許文献１１ではＢｉを添加し、ＡｌＲの式により最終冷延前焼鈍温度を規定している。特許文献１１ではＴｉを含有させ、ＴｉＮを考慮したＡｌＲの式により窒化量を規定している。 In the second technique, Patent Document 9 defines the ratio of N 2 in the atmosphere during finish annealing by the AlR equation. In Patent Document 11, Bi is added, and the annealing temperature before final cold rolling is defined by the AlR equation. In Patent Document 11, Ti is contained, and the amount of nitriding is defined by the AlR equation considering TiN.
第三の技術の場合、一次再結晶粒径の一次再結晶焼鈍温度依存性は小さいが、インヒビター成分、特にＡｌ、Ｎ、さらにはＡｌＮ形成に影響を与えるＴｉの含有量が変動すると、二次再結晶性が不安定になる場合がある。 In the case of the third technique, the primary recrystallization annealing temperature dependence on the primary recrystallization grain size is small, but if the content of Ti that affects the formation of the inhibitor component, particularly Al, N, and further AlN, varies, Recrystallization may become unstable.
ＡｌＲが大きい場合、磁気特性を確保するためには後工程における窒化量を多くする必要がある。この原因は現在次のように考えられている。ＡｌＲが大きいと、最終冷間圧延前焼鈍後にＡｌＮが多く析出し一次粒径が小さくなるが、一次インヒビターの二次インヒビター効果が強くなるので、二次再結晶開始温度は高くなる。そのままでは、高温化に対してインヒビター強度は質的に充分でなく粒径とインヒビターのバランスが崩れて二次再結晶不良となる。高くなった二次再結晶温度に相当すべく窒化により二次インヒビターを強める必要があり、窒化量を増やす必要が生じる。即ち、二次再結晶温度が上がるとインヒビター強度の変化程度は大きくなるため粗大なインヒビターが必要になると考えられる。しかしながら窒化量を大きくすると、グラス皮膜が著しく劣化する。 When AlR is large, it is necessary to increase the amount of nitriding in the subsequent process in order to ensure magnetic properties. The cause is currently considered as follows. When the AlR is large, a large amount of AlN precipitates after the annealing before the final cold rolling and the primary particle size becomes small, but the secondary inhibitor effect of the primary inhibitor becomes strong, so the secondary recrystallization start temperature becomes high. As it is, the inhibitor strength is not qualitatively sufficient for high temperature, and the balance between the particle size and the inhibitor is lost, resulting in a secondary recrystallization failure. In order to correspond to the increased secondary recrystallization temperature, it is necessary to strengthen the secondary inhibitor by nitriding, and it is necessary to increase the amount of nitriding. That is, it is considered that a coarse inhibitor is required because the degree of change in the inhibitor strength increases as the secondary recrystallization temperature increases. However, when the nitriding amount is increased, the glass film is remarkably deteriorated.
他方、ＡｌＲが小さいと最終冷間圧延前焼鈍後にＡｌＮは少なく析出し、一次粒径は大きくなるが、一次インヒビターの二次インヒビター効果が小さいので、二次再結晶開始温度は高くならず、窒化量は少なくて済むが、ＡｌＲが小さ過ぎると、非特許文献１に記載のように、二次再結晶核発生位置が板厚全体に広がるため、表層近傍の先鋭なＧｏｓｓ方位ばかりでなく中心層の粒も二次再結晶し、磁気特性が劣化する。 On the other hand, if the AlR is small, less AlN precipitates after annealing before the final cold rolling, and the primary particle size increases, but the secondary inhibitor effect of the primary inhibitor is small, so the secondary recrystallization start temperature does not increase, and nitriding If the amount of AlR is too small, the secondary recrystallization nucleus generation position spreads over the entire plate thickness as described in Non-Patent Document 1, so that not only the sharp Goss orientation near the surface layer but also the central layer These grains are also secondarily recrystallized and the magnetic properties are deteriorated.
この様に、ＡｌＲが変化すると二次再結晶性、ひいてはＧｏｓｓ方位の先鋭性が変化する。しかしながら溶製段階でＡｌ，Ｎ、Ｔｉの成分範囲を狭い範囲に制御することは困難であるため、これら成分変動の影響を緩和する方策が切望されていた。 Thus, when AlR changes, the secondary recrystallization property, and hence the sharpness of Goss orientation, changes. However, since it is difficult to control the Al, N, and Ti component ranges to a narrow range at the melting stage, there has been a strong demand for measures to alleviate the effects of these component variations.
一方向性電磁鋼板は熱間圧延後多くの工程を経て生産されることは良く知られていることであるが、本発明では、スラブ加熱温度を極端に高くも低くもせず、通常の熱間圧延機で生産でき、また特別なスラブ加熱装置を必要としないことを前提に、成分が不可避的に変動しても熱間圧延後の工程でインヒビター強度を一定に保ち、極めて磁気特性が良好である一方向性電磁鋼板を製造することができる。 It is well known that unidirectional electrical steel sheets are produced through a number of processes after hot rolling, but in the present invention, the slab heating temperature is neither extremely high nor low. On the premise that it can be produced by a rolling mill and does not require a special slab heating device, the inhibitor strength is kept constant in the process after hot rolling even if the components inevitably fluctuate, and the magnetic properties are extremely good. A certain unidirectional electrical steel sheet can be manufactured.
本発明は以下の構成からなる。
（１）質量％で、
Ｃ：０．０３５〜０．０９％、
Ｓｉ：２．５〜４．０％、
酸可溶性Ａｌ（以下、ｓＡｌ）：０．０１８〜０．０３３％、
Ｎ：０．００３〜０．００７％、
Ｓ及び／またはＳｅをＳｅｑ＝Ｓ＋０．４０６Ｓｅで０．００５〜０．０１８％、
Ｍｎ：０．０３〜０．０９％、
Ｔｉを０．００５％以下、
残部がＦｅ及び不可避的不純物からなる一方向性電磁鋼板のスラブを１２００℃以上１３５０℃未満の温度でスラブ加熱を行い、熱間圧延し、熱延板焼鈍と１回の冷間圧延、または中間焼鈍をはさむ２回以上の冷間圧延を行い、熱延板焼鈍もしくは最後の中間焼鈍の後の冷却を７００℃から３００℃までを１０℃／秒以上の速度で冷却を行い、脱炭焼鈍し、焼鈍分離剤を塗布し、脱炭焼鈍後仕上げ焼鈍の二次再結晶開始までの間に鋼板に窒化処理を施し、前記脱炭焼鈍完了後の一次再結晶粒の平均粒径を７μｍ以上１８μｍ未満とし、前記脱炭焼鈍後に、ストリップ走行状態下で水素、窒素、アンモニアの混合ガス中で窒化処理を行い、鋼板窒化後の窒素含有量（ｔＮ）（単位ｐｐｍ）を、０．２５×ｓＡｌ＋９０≦ｔN≦３００ の範囲内とし、最後の冷間圧延の圧下率を８０％以上、９５％以下とし、仕上げ焼鈍を施す一方向性電磁鋼板の製造方法において、熱間圧延後の鋼帯のＮのうちＡｌＮとしての析出率を３０％以下とし、ｓＡｌ、Ｎ、Ｔｉ含有量によって式１で規定される指標ＡｌＲによって、最終冷間圧延前焼鈍の最高温度Ｔ１（℃）を式２の範囲とするとともに、前記熱延板焼鈍もしくは最後の中間焼鈍の焼鈍温度を式２のＴ１（℃）の範囲で、２０〜３６０秒間の一段階焼鈍、又は一段目Ｔ１（℃）で５〜１２０秒間、２段目Ｔ２（℃）８５０〜１０００℃の範囲で１０〜２４０秒間の２段階焼鈍することを特徴とする一方向性電磁鋼板の製造方法。
AlR(ppm)＝sAl(ppm)−27/14×N(ppm)＋27/47.9×Ti(ppm) ・・・式１
−4/3×AlR＋3850/3≦T1≦−4/3×AlR＋4210/3 ・・・式２
（２）スラブの成分が質量％でさらにＳｂとＳｎを、Ｓｂ＋Ｓｎで０．０２〜０．３０％含有することを特徴とする（１）に記載の一方向性電磁鋼板の製造方法。
（３）スラブの成分が、質量％で、さらにＣｕを０．０２〜０．３０％含有することを特徴とする（１）または（２）に記載の一方向性電磁鋼板の製造方法。
The present invention has the following configuration.
(1) In mass%,
C: 0.035 to 0.09%,
Si: 2.5-4.0%
Acid-soluble Al (hereinafter referred to as sAl): 0.018 to 0.033%,
N: 0.003 to 0.007%,
S and / or Se with Seq = S + 0.406Se, 0.005 to 0.018%,
Mn: 0.03 to 0.09%,
Ti is 0.005% or less,
A slab of a unidirectional electrical steel sheet, the balance of which is Fe and inevitable impurities, is slab heated at a temperature of 1200 ° C. or more and less than 1350 ° C., hot-rolled, hot-rolled sheet annealing and one cold rolling, or intermediate Cold rolling is performed at least twice with annealing, and cooling after hot-rolled sheet annealing or the last intermediate annealing is performed from 700 ° C to 300 ° C at a rate of 10 ° C / second or more and decarburized annealing. The steel sheet is subjected to nitriding treatment before the secondary recrystallization of the annealing after the decarburization annealing and after the decarburization annealing, and the average grain size of the primary recrystallized grains after the completion of the decarburization annealing is not less than 7 μm and not more than 18 μm. After decarburization annealing, nitriding treatment was performed in a mixed gas of hydrogen, nitrogen and ammonia under strip running conditions, and the nitrogen content (tN) (unit: ppm) after nitriding the steel sheet was 0.25 × sAl + 90 Within the range of ≦ tN ≦ 300 Final rolling reduction of cold rolling of 80% or more, and 95% or less, in the manufacturing method of the grain-oriented electrical steel sheet subjected to finish annealing, the precipitation rate of the AlN of N in the steel strip after hot rolling 30 % and less, SAL, N, by the index AlR defined by equation 1 by the Ti content, the final cold rolling before annealing at a maximum temperature T1 (° C.) with the range of formula 2, wherein hot rolled sheet annealing or The annealing temperature of the final intermediate annealing is in the range of T1 (° C) in Formula 2 for one-step annealing for 20 to 360 seconds, or for the first step T1 (° C) for 5 to 120 seconds, and for the second step T2 (° C) 850 A method for producing a unidirectional electrical steel sheet, characterized by performing two-stage annealing for 10 to 240 seconds in a range of 1000 ° C.
AlR (ppm) = sAl (ppm) −27 / 14 × N (ppm) + 27 / 47.9 × Ti (ppm) ・ ・ ・ Equation 1
−4 / 3 × AlR + 3850/3 ≦ T1 ≦ −4 / 3 × AlR + 4210/3 ・ ・ ・ Equation 2
( 2 ) The method for producing a unidirectional electrical steel sheet according to (1), wherein the slab component is contained by mass%, and Sb and Sn are further contained by 0.02 to 0.30% as Sb + Sn.
( 3 ) The method for producing a unidirectional electrical steel sheet according to (1) or ( 2 ), wherein the slab component contains 0.02 to 0.30% by mass and further Cu.
本発明においては、従来の一方向性電磁鋼板の熱延加熱時の超高温度を脱却すると共に低温加熱の弊害を取り除いて磁気特性の優れる一方向性電磁鋼板が製造可能になる。 In the present invention, it becomes possible to manufacture a unidirectional electrical steel sheet having excellent magnetic properties by removing the ultra-high temperature during hot rolling of a conventional unidirectional electrical steel sheet and removing the adverse effects of low temperature heating.
以下に本発明を詳細に説明する。発明者らは、第三の技術のうち完全固溶窒化型の製造方法について鋭意検討を行った結果、熱間圧延で固溶・凍結されたインヒビター物質の形態を最終冷間圧延直前の焼鈍条件で変化させることができることを見出した。特に、そのときのパラメーターとしてＡｌＲが有効であり、その焼鈍の温度により二次再結晶の安定性が大きく影響されることを見出したのである。 The present invention is described in detail below. As a result of intensive studies on the manufacturing method of the complete solid solution nitriding type in the third technique, the inventors have determined that the form of the inhibitor substance dissolved and frozen by hot rolling is the annealing condition immediately before the final cold rolling. I found that it can be changed. In particular, it was found that AlR is effective as a parameter at that time, and the stability of secondary recrystallization is greatly influenced by the annealing temperature.
次に本発明におけるスラブの成分範囲の限定理由について述べる。 Next, the reason for limiting the component range of the slab in the present invention will be described.
Ｃは、０．０３５％より少ないと一次再結晶集合組織が適切でなくなり、０．０９％を超えると脱炭が困難になり工業生産に適していない。 When C is less than 0.035%, the primary recrystallization texture becomes unsuitable, and when it exceeds 0.09%, decarburization becomes difficult and is not suitable for industrial production.
Ｓｉは、２．５％より少ないと良好な鉄損が得られず、４．０％を超えると冷延が極めて困難となり工業生産に適していない。 If Si is less than 2.5%, good iron loss cannot be obtained, and if it exceeds 4.0%, cold rolling becomes extremely difficult and is not suitable for industrial production.
ＡｌはＮと結合してＡｌＮを形成し、主に一次・二次インヒビターとして機能する。溶製段階で含有され熱間圧延で凍結（強制的に固溶状態から冷却され熱延鋼帯で析出していない）され最終冷間圧延直前の焼鈍で析出するものは、主に一次インヒビターとして機能するが二次インヒビターとしても機能する。さらに、脱炭焼鈍後二次再結晶開始までで窒化され形成されたＡｌＮは二次インヒビターとして機能する。即ち、このＡｌＮは、窒化前に形成されるものと窒化後高温焼鈍時に形成されるものがあり、この両方のＡｌＮ量確保のために酸可溶性Ａｌは、０．０１８〜０．０３３％必要である。この範囲を外れると、たとえ本発明の技術を用いても、低い場合は、二次インヒビターとしての働きが不充分な為良好なＧｏｓｓ方位を持った二次再結晶粒を安定的に得られず、高い場合には、後工程の必要窒化量が増大し、グラス皮膜に甚大なダメージを与える。 Al combines with N to form AlN, and mainly functions as a primary and secondary inhibitor. Contained in the melting stage, frozen by hot rolling (forced to cool from the solid solution state and not precipitated in the hot-rolled steel strip), and precipitated by annealing immediately before the final cold rolling, mainly as a primary inhibitor It functions but also functions as a secondary inhibitor. Further, AlN formed by nitriding until the start of secondary recrystallization after decarburization annealing functions as a secondary inhibitor. That is, this AlN is formed before nitriding and formed during high-temperature annealing after nitriding. In order to secure the amount of both AlN, 0.018 to 0.033% of acid-soluble Al is required. is there. Outside this range, even if the technique of the present invention is used, if it is low, secondary recrystallized grains having good Goss orientation cannot be stably obtained because the function as a secondary inhibitor is insufficient. If it is high, the amount of nitridation required in the subsequent process increases, and the glass film is seriously damaged.
Ｎは、０．００７％を越えると、実際の工業生産では熱延のスラブ加熱温度を１３５０℃超にする必要が生じる。より好ましくは、０．００６％以下である。０．００３％未満では、安定した一次インヒビター効果が得られず二次再結晶不良となる。 If N exceeds 0.007%, in actual industrial production, the slab heating temperature for hot rolling needs to be higher than 1350 ° C. More preferably, it is 0.006% or less. If it is less than 0.003%, a stable primary inhibitor effect cannot be obtained, resulting in poor secondary recrystallization.
ＳおよびＳｅはＭｎ，Ｃｕと結合して、主に一次インヒビターとして作用する。Ｓｅｑ＝Ｓ＋０．４０６×Ｓｅが０．０１８％を超えると、熱間圧延時のスラブ加熱温度が高くなる。また、０．００５％未満とすると、一次インヒビターとしての効果が弱くなる。望ましくは、０．０１０〜０．０１５％である。 S and Se combine with Mn and Cu and act mainly as a primary inhibitor. When Seq = S + 0.406 × Se exceeds 0.018%, the slab heating temperature during hot rolling increases. Moreover, when less than 0.005%, the effect as a primary inhibitor will become weak. Desirably, it is 0.010 to 0.015%.
Ｍｎは０．０３％より少ないと熱間圧延後に割れが発生しやすく、歩留まりが低下する。一方０．０９％を超えるとＭｎＳ、ＭｎＳｅが多くなりすぎ熱間圧延時のスラブ加熱温度を高くせねばならなくなる。望ましくは、０．０４〜０．０６％である。 If Mn is less than 0.03%, cracks are likely to occur after hot rolling, and the yield decreases. On the other hand, if it exceeds 0.09%, MnS and MnSe will increase so much that the slab heating temperature during hot rolling must be increased. Desirably, it is 0.04 to 0.06%.
Ｃｕは、スラブを１２００℃以上で加熱する本発明の条件で熱延すると、ＳやＳｅとともに微細な析出物を形成し、一次インヒビター効果を発揮する。また、この析出物はＡｌＮの分散をより均一にする析出核ともなり二次インヒビターの役割も演じ、この効果が二次再結晶を良好ならしめる。０．０２％より少ないと上記効果が減じ安定生産が難しくなり、０．３０％を超えると上記効果が飽和するとともに、熱延時に「カッパーヘゲ」なる表面疵の原因になる。 When hot rolling is performed under the conditions of the present invention in which the slab is heated at 1200 ° C. or higher, Cu forms fine precipitates together with S and Se and exhibits a primary inhibitor effect. The precipitates also serve as precipitation nuclei that make the dispersion of AlN more uniform and also play a role of secondary inhibitors, and this effect makes secondary recrystallization good. If the content is less than 0.02%, the above effect is reduced and stable production becomes difficult. If the content exceeds 0.30%, the above effect is saturated, and it causes surface flaws such as “copper lashes” during hot rolling.
ＳｎとＳｂは、良く知られている粒界偏析元素である。本発明はＡｌを含有しているため、仕上げ焼鈍の条件によっては焼鈍分離剤から放出される水分によりＡｌが酸化されてコイル位置でインヒビター強度が変動し、磁気特性がコイル位置で変動する場合がある。この対策として仕上げ焼鈍における緻密なグラス皮膜の早期形成の他、粒界偏析元素の添加により酸化を防止する方法がある。Ｓｎ＋Ｓｂで０．０２％未満であるとこの効果が極めて小さい。一方０．３０％を超えると脱炭焼鈍時に酸化されにくくグラス皮膜形成が不十分となり、脱炭焼鈍性を著しく阻害する。 Sn and Sb are well-known grain boundary segregation elements. Since the present invention contains Al, depending on the conditions of finish annealing, Al is oxidized by moisture released from the annealing separator, and the inhibitor strength varies at the coil position, and the magnetic characteristics may vary at the coil position. is there. As a countermeasure, there is a method of preventing oxidation by adding a grain boundary segregation element in addition to the early formation of a dense glass film in finish annealing. If Sn + Sb is less than 0.02%, this effect is extremely small. On the other hand, if it exceeds 0.30%, it is difficult to be oxidized at the time of decarburization annealing, and the glass film formation becomes insufficient, and the decarburization annealing property is remarkably inhibited.
Ｔｉは０．００５％を超えると二次再結晶が不安定になり、安定化するためには窒化後窒素含有量を０．０３０％以上とする必要が生じグラス皮膜不良の原因となる。 If Ti exceeds 0.005%, secondary recrystallization becomes unstable, and in order to stabilize, it is necessary to make the nitrogen content after nitriding 0.030% or more, which causes a glass film defect.
その他、一方向性電磁鋼板の諸特性を向上させる周知の元素を添加できる。その好適添加範囲は、Ｐ及びＣｒ：０．０２〜０．３０％、Ｍｏ，Ｃｄ：０．００８〜０．３％、そしてＮｉ：０．０３〜０．３０％であり、これらの各元素についても、単独使用及び複合使用いずれもが可能である。 In addition, known elements that improve various characteristics of the unidirectional electrical steel sheet can be added. The preferred addition ranges are P and Cr: 0.02 to 0.30%, Mo, Cd: 0.008 to 0.3%, and Ni: 0.03 to 0.30%. Both can be used alone or in combination.
次に本発明における製造工程の限定理由について述べる。 Next, the reasons for limiting the manufacturing process in the present invention will be described.
本発明の方法では、第一に、公知の連続鋳造法により初期の厚みが１５０ｍｍから３５０ｍｍの範囲、好ましくは２２０ｍｍから２８０ｍｍの範囲のスラブを製造する。この代わりに、スラブは初期の厚みが約３０ｍｍから７０ｍｍの範囲のいわゆる薄いスラブであってもよく、この場合は、熱延鋼帯を製造する際、中間厚みに粗加工をする必要がないとの利点がある。 In the method of the present invention, first, a slab having an initial thickness in the range of 150 mm to 350 mm, preferably in the range of 220 mm to 280 mm, is manufactured by a known continuous casting method. Alternatively, the slab may be a so-called thin slab having an initial thickness in the range of about 30 mm to 70 mm, and in this case, when manufacturing a hot-rolled steel strip, there is no need to roughen the intermediate thickness. There are advantages.
スラブ加熱温度を１２００〜１３５０℃に限定した理由は、１２００℃未満とすることは前記第二の技術に相当し、良好なＧｏｓｓ方位の二次再結晶を達成するためのインヒビター物質の固溶が熱力学的に不可能である。またスラブ加熱温度が低いと熱間圧延機の圧延負荷が大きくなりミスロールの原因や平坦度（クラウン）不良が起き易い。一方１３５０℃を超えることは前記第一の技術に相当し、従来から良く知られている工業生産上の困難性が生じる。望ましい範囲は１２８０℃〜１３３０℃である。 The reason why the slab heating temperature is limited to 1200 to 1350 ° C. is that it is less than 1200 ° C., which corresponds to the second technique described above, and the solid solution of the inhibitor substance for achieving good Goss orientation secondary recrystallization Thermodynamically impossible. In addition, when the slab heating temperature is low, the rolling load of the hot rolling mill is increased, and the cause of misroll and flatness (crown) failure are likely to occur. On the other hand, exceeding 1350 ° C. corresponds to the first technique, which causes a well-known difficulty in industrial production. A desirable range is 1280 ° C to 1330 ° C.
スラブ加熱方法として通常はガス加熱方法が用いられるが、誘導加熱、直接通電加熱を用いると均一に加熱する点で望ましく、これらの特別な加熱方法において形状を確保するため、分塊圧延を鋳込みスラブに施しても何ら問題ない。また、加熱温度を１３００℃以上とする場合は、この分塊圧延により集合組織の改善を施しＣ量を減じてもよい。 A gas heating method is usually used as the slab heating method. However, induction heating and direct current heating are desirable in terms of uniform heating, and in order to ensure the shape in these special heating methods, the partial slab is cast into a slab. There is no problem even if given to. Moreover, when heating temperature shall be 1300 degreeC or more, the texture may be improved by this partial rolling and the amount of C may be reduced.
本発明においては、熱間圧延後の段階においてインヒビター物質が凍結されていることが必要である。この指標としては、ＮのＡｌＮとしての析出率が３０％以下である。３０％を超えると熱間圧延鋼帯内でのインヒビター物質形態が位置的不均一となり工業的に安定性生産できない。 In the present invention, it is necessary that the inhibitor substance is frozen at the stage after hot rolling. As this index, the precipitation rate of N as AlN is 30% or less. If it exceeds 30%, the form of the inhibitor substance in the hot-rolled steel strip becomes non-uniform in position, and industrially stable production cannot be performed.
最終冷間圧延前の焼鈍、すなわち熱延板焼鈍もしくは最後の中間焼鈍は、熱間圧延時に凍結された鋼帯内の組織の均質化とインヒビターの微細分散にとって重要である。即ち、最終冷間圧延前に熱延時履歴差による不均一性・インヒビター物質の微細析出させるために１回以上の連続焼鈍を行うことが必須である。この焼鈍条件のうち特に最高温度が非常に重要で、これによりＡｌＲの変動による二次再結晶性の変動が吸収されることを見出した。即ち、最高温度をＴ１（℃）とすると満たすべき条件は、式２の範囲である。
AlR(ppm)＝sAl(ppm)−27/14×N(ppm)＋27/47.9×Ti(ppm) ・・・式１
−4/3×AlR＋3850/3≦T1≦−4/3×AlR＋4210/3 ・・・式２
これより低いとＧｏｓｓの先鋭性が低下し、高いと二次再結晶が不安定となり窒化量を著しく多くすることが求められ工業生産的に現実的でない。
Annealing before the final cold rolling, that is, hot-rolled sheet annealing or the final intermediate annealing is important for the homogenization of the structure in the steel strip frozen during the hot rolling and the fine dispersion of the inhibitors. That is, it is indispensable to perform one or more continuous annealings before the final cold rolling in order to cause non-uniformity due to the difference in hot rolling history and fine precipitation of the inhibitor substance. Of these annealing conditions, the maximum temperature is particularly important, and it has been found that variations in secondary recrystallization due to variations in AlR are absorbed. That is, the condition to be satisfied when the maximum temperature is T1 (° C.) is in the range of Equation 2.
AlR (ppm) = sAl (ppm) −27 / 14 × N (ppm) + 27 / 47.9 × Ti (ppm) ・ ・ ・ Equation 1
−4 / 3 × AlR + 3850/3 ≦ T1 ≦ −4 / 3 × AlR + 4210/3 ・ ・ ・ Equation 2
If it is lower than this, the sharpness of Goss decreases, and if it is higher, secondary recrystallization becomes unstable and it is required to significantly increase the amount of nitriding, which is not practical for industrial production.
また、この場合の焼鈍条件のヒートパターンとしては２種類ある。一つは２段階による方法で、その１段目は温度を前記式２に示すＴ１（℃）の範囲で５〜１２０秒間とし、２段目の温度Ｔ２（℃）を８５０〜１０００℃の範囲で１０秒から２４０秒間とする方法である。もう一つは一段サイクルで、その温度を前記式２に示すＴ１（℃）の範囲で２０〜３６０秒間とする方法である。２つの方法では冷却条件が異なるため、集合組織に及ぼす影響が生じるが、本願発明が課題とするインヒビター状態の調整はＴ１によって決まるので、いずれの方法を用いるかは、現状の設備制約などにより適宜決めることが出来る。 Moreover, there are two types of heat patterns for the annealing conditions in this case. One is a two-stage method, and the first stage has a temperature in the range of T1 (° C) shown in Equation 2 for 5 to 120 seconds, and the second stage temperature T2 (° C) is in the range of 850 to 1000 ° C. This is a method of 10 to 240 seconds. The other is a one-stage cycle in which the temperature is set to 20 to 360 seconds in the range of T1 (° C.) shown in Formula 2 above. Since the two methods have different cooling conditions, there is an effect on the texture. However, since the adjustment of the inhibitor state, which is the subject of the present invention, is determined by T1, which method is used depends on the current equipment constraints and the like. I can decide.
以上の焼鈍温度からの冷却は、インヒビターを固定するため重要である。古くは特公昭４６−２３８２０号公報に７５０〜９５０℃から４００℃までを２〜２００秒で行うことが示されている。本発明では、７００℃以上から３００℃以下までの冷却が１０℃／秒以上が確保されれば充分である。 Cooling from the above annealing temperature is important for fixing the inhibitor. In the past, Japanese Patent Publication No. 46-23820 discloses that the temperature from 750 to 950 ° C. to 400 ° C. is carried out in 2 to 200 seconds. In the present invention, it is sufficient if cooling from 700 ° C. to 300 ° C. is ensured at 10 ° C./second or more.
冷間圧延は１回もしくは中間焼鈍をはさむ２回以上で行われる。このとき最終の冷間圧延における最終冷延率が８０％未満であると、一次再結晶集合組織中のＧｏｓｓ方位粒の方位集積度が得難いため高磁束密度が確保し難くなる。一方９５％を超えると一次再結晶集合組織中Ｇｏｓｓ方位粒数が極端に少なくなるため二次再結晶が不安定になる。 Cold rolling is performed once or twice or more with intermediate annealing. At this time, if the final cold rolling rate in the final cold rolling is less than 80%, it is difficult to obtain the orientation accumulation degree of Goss orientation grains in the primary recrystallization texture, so that it is difficult to secure a high magnetic flux density. On the other hand, if it exceeds 95%, the number of Goss orientation grains in the primary recrystallization texture becomes extremely small, and secondary recrystallization becomes unstable.
最終冷間圧延は常温で実施してもよいが、少なくとも１パスを１００〜３００℃の温度範囲に１分以上保つと一次再結晶集合組織が改善され磁気特性が極めて良好になる。 The final cold rolling may be carried out at room temperature, but if at least one pass is kept in the temperature range of 100 to 300 ° C. for 1 minute or longer, the primary recrystallization texture is improved and the magnetic properties become extremely good.
脱炭焼鈍は周知の方法で行われる。このとき室温から６００〜８００℃までの加熱速度を１００℃／秒以上の急速加熱とすると、一次再結晶集合組織が改善され磁気特性が良好になる。加熱速度を確保するためには種々の方法が考えられる。即ち、抵抗加熱、誘導加熱、直接エネルギー付与加熱等である。加熱速度を早くすると一次再結晶集合組織においてＧｏｓｓ方位粒が多くなり二次再結晶粒径が小さくなることは特開平６−５１８８７号公報等に開示されている。本発明では、前記加熱速度が１００℃／秒でも効果があり、望ましくは１２０℃／秒以上である。急速加熱を行う温度範囲を６００〜８００℃としたのは、これ未満では再結晶が完了せず効果が少ないためであり、一方この範囲を超えて急速加熱を継続しても、一次再結晶での集合組織に大きな影響を与えるＧｏｓｓ核の生成は８００℃までで完了するため、効果は小さい。 Decarburization annealing is performed by a known method. At this time, when the heating rate from room temperature to 600 to 800 ° C. is rapid heating of 100 ° C./second or more, the primary recrystallization texture is improved and the magnetic properties are improved. Various methods are conceivable for securing the heating rate. That is, resistance heating, induction heating, direct energy application heating, and the like. It is disclosed in JP-A-6-51887 and the like that when the heating rate is increased, the number of Goss orientation grains increases in the primary recrystallization texture and the secondary recrystallization grain size decreases. In the present invention, the heating rate is effective even at 100 ° C./second, desirably 120 ° C./second or more. The reason why the temperature range for rapid heating is set to 600 to 800 ° C. is that recrystallization is not completed when the temperature is lower than this, and the effect is small. On the other hand, even if rapid heating is continued beyond this range, primary recrystallization is not possible. Since the formation of Goss nuclei having a great influence on the texture of this material is completed up to 800 ° C., the effect is small.
脱炭焼鈍完了後の一次再結晶粒の平均粒径は、第二の技術においては特開平７−２５２５３２号公報等に、一次再結晶粒の平均粒径を１８〜３５μｍとすることが開示されているが、第三の技術に関する本発明では、一次再結晶粒の平均粒径を７μｍ以上１８μｍ未満とすることで、磁気特性（特に鉄損）を更に良好ならしめることができる。即ち、粒径が小さければ、単位体積内に存在する一次再結晶粒数が増えることを意味する。更に、一次再結晶粒径が小さい場合、粒成長の観点から、一次再結晶の段階で二次再結晶の核となるＧｏｓｓ方位粒の体積分率が多くなり（Materials Science Forum Vol.204-206,Part2:pp:631）、結果としてＧｏｓｓ方位粒の絶対数は、例えば一次再結晶粒の平均粒径が１８〜３５μｍの場合と比べると、５倍程度も多くなりＧｏｓｓ粒の選択成長性が著しく向上し、磁気特性が向上する。 In the second technique, the average grain size of primary recrystallized grains after completion of decarburization annealing is disclosed in Japanese Patent Laid-Open No. 7-252532, etc., so that the average grain size of primary recrystallized grains is 18 to 35 μm. However, in the present invention relating to the third technique, the magnetic properties (particularly iron loss) can be further improved by setting the average primary recrystallized grain size to 7 μm or more and less than 18 μm. That is, if the particle size is small, it means that the number of primary recrystallized grains existing in the unit volume increases. Furthermore, when the primary recrystallized grain size is small, from the viewpoint of grain growth, the volume fraction of Goss orientation grains that become the nucleus of secondary recrystallization at the primary recrystallization stage increases (Materials Science Forum Vol.204-206). Part 2: pp: 631) As a result, the absolute number of Goss orientation grains is about 5 times larger than that of the primary recrystallized grains having an average grain size of 18 to 35 μm, for example. Significantly improved and magnetic properties are improved.
また、一次再結晶粒の平均粒径が小さいと、二次再結晶の駆動力が大きくなり、仕上げ焼鈍中昇温段階のより低温段階で二次再結晶を開始させることができる。仕上げ焼鈍をコイル状で行っている現状では、昇温中の高温域ほどコイル内外の温度差が広がるので、上述の二次再結晶温度の低温化によってコイル部位による二次再結晶の不均一性が著しく減少し、磁気特性が極めて安定する。 In addition, when the average particle size of the primary recrystallized grains is small, the driving force for secondary recrystallization is increased, and secondary recrystallization can be started at a lower temperature stage than the temperature raising stage during finish annealing. In the present situation where finish annealing is performed in a coil shape, the temperature difference between the inside and outside of the coil increases as the temperature rises, so the secondary recrystallization due to the coil site is not uniform due to the lowering of the secondary recrystallization temperature described above. Is significantly reduced and the magnetic properties are extremely stable.
但し、一次再結晶粒の平均粒径が７μｍ未満の場合、粒成長駆動力が大きくなりすぎて、二次再結晶粒方位のＧｏｓｓ方位からの分散が大きくなり、磁束密度の低下をまねくので好ましくない。 However, when the average grain size of the primary recrystallized grains is less than 7 μm, the grain growth driving force becomes too large, and the dispersion of the secondary recrystallized grain orientation from the Goss orientation becomes large, which leads to a decrease in magnetic flux density. Absent.
脱炭焼鈍後二次再結晶開始前に鋼板に窒化処理を施すことは本発明では必須である。その方法は、仕上げ焼鈍時の焼鈍分離剤に窒化物（ＣｒＮ，ＭｎＮ等）を混合させる方法や、脱炭焼鈍後にストリップを走行させた状態下でアンモニアを含んだ雰囲気で窒化させる方法がある。どちらの方法を採用しても良いが、後者の方が工業的に安定している。 It is essential in the present invention that the steel sheet is subjected to nitriding after decarburization annealing and before the start of secondary recrystallization. As the method, there are a method of mixing a nitride (CrN, MnN, etc.) with an annealing separator at the time of final annealing, and a method of nitriding in an atmosphere containing ammonia in a state where the strip is run after decarburization annealing. Either method may be adopted, but the latter is more industrially stable.
このときの鋼板窒化後の窒素含有量（ｔＮ）（単位ppm）は、ｓＡｌの含有量に応じて、０．２５×ｓＡｌ＋９０≦ｔN≦３００ の範囲とすることが好ましい。これより少ないと二次再結晶が不安定となり、一方窒化後の総窒素含有量が０．０３０％を超えると、過剰なＮが仕上げ焼鈍中に吹き出してグラス皮膜を破壊する場合がある。望ましくは、窒素増量として０．０１８％〜０．０２２％である。 At this time, the nitrogen content (tN) (unit: ppm) after nitriding the steel sheet is preferably in the range of 0.25 × sAl + 90 ≦ tN ≦ 300 according to the content of sAl. If it is less than this, secondary recrystallization becomes unstable, while if the total nitrogen content after nitriding exceeds 0.030%, excessive N may blow out during finish annealing and destroy the glass film. Desirably, the nitrogen increase is 0.018% to 0.022%.
表２に示す成分を有するスラブを１３００〜１３３０℃で加熱後、できるだけ高温で熱延を完了し２．３ｍｍ厚の鋼帯を急冷して５００〜６００℃で巻き取った。そして、その熱延鋼帯を最高温度（Ｔ１）を１０５０、１０８０、１１１０、１１４０、及び１１７０℃として１５０秒間窒素雰囲気中で均熱し、冷却しその後７３０℃まで空冷後１００℃の熱湯に焼き入れた。その後、１５０℃〜２１０℃で３パス確保して冷間圧延で０．２８５ｍｍとした。その後、８５０℃で１５０秒の均熱の脱炭・一次再結晶焼鈍を水素７５％窒素２５％の露点７０℃の雰囲気で行った。これに引き続いて窒素含有量を０．０１８〜０．０２１％となるように走行するストリップ状でアンモニアと窒素の混合ガス中で窒化した。そして、ＭｇＯを主成分とする焼鈍分離剤を塗布した。次に、窒素２５％、水素７５％の雰囲気で１２００℃まで１５℃／ｈで加熱し１２００℃に到達後は、水素１００％雰囲気で純化処理を行い仕上げ焼鈍を行った。その後通常用いられる燐酸アルミニウムを主成分とする張力絶縁コーティングを塗布した後、磁気特性を測定した。 After heating the slab which has a component shown in Table 2 at 1300-1330 degreeC, hot rolling was completed as high as possible, the 2.3 mm-thick steel strip was quenched, and it wound up at 500-600 degreeC. The hot-rolled steel strip is heated at a maximum temperature (T1) of 1050, 1080, 1110, 1140, and 1170 ° C. in a nitrogen atmosphere for 150 seconds, cooled, air-cooled to 730 ° C., and then quenched in hot water at 100 ° C. It was. Thereafter, 3 passes were secured at 150 ° C. to 210 ° C., and the thickness was 0.285 mm by cold rolling. Thereafter, soaking and primary recrystallization annealing at 850 ° C. for 150 seconds was performed in an atmosphere with a dew point of 70 ° C. of 75% hydrogen and 25% nitrogen. This was followed by nitriding in a mixed gas of ammonia and nitrogen in the form of a strip running so that the nitrogen content was 0.018 to 0.021%. And the annealing separation agent which has MgO as a main component was apply | coated. Next, heating was performed at 15 ° C./h up to 1200 ° C. in an atmosphere of 25% nitrogen and 75% hydrogen, and after reaching 1200 ° C., purification treatment was performed in a 100% hydrogen atmosphere and finish annealing was performed. Then, after applying a commonly used tensile insulating coating mainly composed of aluminum phosphate, the magnetic properties were measured.
結果を表３に示す。 The results are shown in Table 3.
表４に示す成分を有するスラブを１３００〜１３３０℃で加熱後、できるだけ高温で熱延を完了し２．１ｍｍ厚の鋼帯を急冷して５００〜６００℃で巻き取った。そして、その熱延鋼帯を最高温度（Ｔ１）を１０５０、１０８０、１１１０、１１４０、及び１１７０℃として１５秒間窒素雰囲気中で均熱し、９００℃までに１５０秒で冷却しその後７５０℃まで空冷後１００℃の熱湯に焼き入れた。その後、１５０℃〜２３０℃で３パス確保して冷間圧延で０．２２０ｍｍとした。その後、８５０℃で９０秒の均熱の脱炭・一次再結晶焼鈍を水素７５％窒素２５％の露点６９℃の雰囲気で行った。これに引き続いて窒素含有量を０．０１８〜０．０２１％となるように走行するストリップ状でアンモニアと窒素の混合ガス中で窒化した。そして、ＭｇＯを主成分とする焼鈍分離剤を塗布した。次に、窒素２５％、水素７５％の雰囲気で１２００℃まで１５℃／ｈで加熱し１２００℃に到達後は、水素１００％雰囲気で純化処理を行い仕上げ焼鈍を行った。その後通常用いられる燐酸アルミニウムを主成分とする張力絶縁コーティングを塗布した後、磁気特性を測定した。 After heating the slab which has a component shown in Table 4 at 1300-1330 degreeC, hot rolling was completed as high as possible, the 2.1 mm-thick steel strip was quenched, and it wound up at 500-600 degreeC. The hot-rolled steel strip was heated at a maximum temperature (T1) of 1050, 1080, 1110, 1140, and 1170 ° C. in a nitrogen atmosphere for 15 seconds, cooled to 900 ° C. in 150 seconds, and then air-cooled to 750 ° C. Quenched in hot water at 100 ° C. Thereafter, three passes were secured at 150 ° C. to 230 ° C., and the thickness was 0.220 mm by cold rolling. Thereafter, soaking and primary recrystallization annealing at 850 ° C. for 90 seconds was performed in an atmosphere having a dew point of 69 ° C. with 75% hydrogen and 25% nitrogen. This was followed by nitriding in a mixed gas of ammonia and nitrogen in the form of a strip running so that the nitrogen content was 0.018 to 0.021%. And the annealing separation agent which has MgO as a main component was apply | coated. Next, heating was performed at 15 ° C./h up to 1200 ° C. in an atmosphere of 25% nitrogen and 75% hydrogen, and after reaching 1200 ° C., purification treatment was performed in a 100% hydrogen atmosphere and finish annealing was performed. Then, after applying a commonly used tensile insulating coating mainly composed of aluminum phosphate, the magnetic properties were measured.
結果を表５に示す。 The results are shown in Table 5.
Claims (3)
Ｃ：０．０３５〜０．０９％、
Ｓｉ：２．５〜４．０％、
酸可溶性Ａｌ（以下、ｓＡｌ）：０．０１８〜０．０３３％、
Ｎ：０．００３〜０．００７％、
Ｓ及び／またはＳｅをＳｅｑ＝Ｓ＋０．４０６Ｓｅで０．００５〜０．０１８％、
Ｍｎ：０．０３〜０．０９％、
Ｔｉを０．００５％以下、
残部がＦｅ及び不可避的不純物からなる一方向性電磁鋼板のスラブを１２００℃以上１３５０℃未満の温度でスラブ加熱を行い、熱間圧延し、熱延板焼鈍と１回の冷間圧延、または中間焼鈍をはさむ２回以上の冷間圧延を行い、熱延板焼鈍もしくは最後の中間焼鈍の後の冷却を７００℃から３００℃までを１０℃／秒以上の速度で冷却を行い、脱炭焼鈍し、焼鈍分離剤を塗布し、脱炭焼鈍後仕上げ焼鈍の二次再結晶開始までの間に鋼板に窒化処理を施し、前記脱炭焼鈍完了後の一次再結晶粒の平均粒径を７μｍ以上１８μｍ未満とし、前記脱炭焼鈍後に、ストリップ走行状態下で水素、窒素、アンモニアの混合ガス中で窒化処理を行い、鋼板窒化後の窒素含有量（ｔＮ）（単位ｐｐｍ）を、０．２５×ｓＡｌ＋９０≦ｔN≦３００ の範囲内とし、最後の冷間圧延の圧下率を８０％以上、９５％以下とし、仕上げ焼鈍を施す一方向性電磁鋼板の製造方法において、熱間圧延後の鋼帯のＮのうちＡｌＮとしての析出率を３０％以下とし、ｓＡｌ、Ｎ、Ｔｉ含有量によって式１で規定される指標ＡｌＲによって、最終冷間圧延前焼鈍の最高温度Ｔ１（℃）を式２の範囲とするとともに、前記熱延板焼鈍もしくは最後の中間焼鈍の焼鈍温度を式２のＴ１（℃）の範囲で、２０〜３６０秒間の一段階焼鈍、又は一段目Ｔ１（℃）で５〜１２０秒間、２段目Ｔ２（℃）８５０〜１０００℃の範囲で１０〜２４０秒間の２段階焼鈍することを特徴とする一方向性電磁鋼板の製造方法。
AlR(ppm)＝sAl(ppm)−27/14×N(ppm)＋27/47.9×Ti(ppm) ・・・式１
−4/3×AlR＋3850/3≦T1≦−4/3×AlR＋4210/3 ・・・式２ % By mass
C: 0.035 to 0.09%,
Si: 2.5-4.0%
Acid-soluble Al (hereinafter referred to as sAl): 0.018 to 0.033%,
N: 0.003 to 0.007%,
S and / or Se with Seq = S + 0.406Se, 0.005 to 0.018%,
Mn: 0.03 to 0.09%,
Ti is 0.005% or less,
A slab of a unidirectional electrical steel sheet, the balance of which is Fe and inevitable impurities, is slab heated at a temperature of 1200 ° C. or more and less than 1350 ° C., hot-rolled, hot-rolled sheet annealing and one cold rolling, or intermediate Cold rolling is performed at least twice with annealing, and cooling after hot-rolled sheet annealing or the last intermediate annealing is performed from 700 ° C to 300 ° C at a rate of 10 ° C / second or more and decarburized annealing. The steel sheet is subjected to nitriding treatment before the secondary recrystallization of the annealing after the decarburization annealing and after the decarburization annealing, and the average grain size of the primary recrystallized grains after the completion of the decarburization annealing is not less than 7 μm and not more than 18 μm. After decarburization annealing, nitriding treatment was performed in a mixed gas of hydrogen, nitrogen and ammonia under strip running conditions, and the nitrogen content (tN) (unit: ppm) after nitriding the steel sheet was 0.25 × sAl + 90 Within the range of ≦ tN ≦ 300 Final rolling reduction of cold rolling of 80% or more, and 95% or less, in the manufacturing method of the grain-oriented electrical steel sheet subjected to finish annealing, the precipitation rate of the AlN of N in the steel strip after hot rolling 30 % and less, SAL, N, by the index AlR defined by equation 1 by the Ti content, the final cold rolling before annealing at a maximum temperature T1 (° C.) with the range of formula 2, wherein hot rolled sheet annealing or The annealing temperature of the final intermediate annealing is in the range of T1 (° C) in Formula 2 for one-step annealing for 20 to 360 seconds, or for the first step T1 (° C) for 5 to 120 seconds, and for the second step T2 (° C) 850 A method for producing a unidirectional electrical steel sheet, characterized by performing two-stage annealing for 10 to 240 seconds in a range of 1000 ° C.
AlR (ppm) = sAl (ppm) −27 / 14 × N (ppm) + 27 / 47.9 × Ti (ppm) ・ ・ ・ Equation 1
−4 / 3 × AlR + 3850/3 ≦ T1 ≦ −4 / 3 × AlR + 4210/3 ・ ・ ・ Equation 2
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| JP4823719B2 (en) * | 2006-03-07 | 2011-11-24 | 新日本製鐵株式会社 | Method for producing grain-oriented electrical steel sheet with extremely excellent magnetic properties |
| KR101062127B1 (en) | 2006-05-24 | 2011-09-02 | 신닛뽄세이테쯔 카부시키카이샤 | Method for manufacturing directional electromagnetic steel sheet with high magnetic flux density |
| KR101070064B1 (en) | 2006-05-24 | 2011-10-04 | 신닛뽄세이테쯔 카부시키카이샤 | Process for producing grain-oriented magnetic steel sheet with high magnetic flux density |
| CN101545072B (en) * | 2008-03-25 | 2012-07-04 | 宝山钢铁股份有限公司 | Method for producing oriented silicon steel having high electromagnetic performance |
| JP5428188B2 (en) * | 2008-04-15 | 2014-02-26 | 新日鐵住金株式会社 | Method for producing grain-oriented electrical steel sheet |
| JP4800442B2 (en) * | 2008-09-10 | 2011-10-26 | 新日本製鐵株式会社 | Method for producing grain-oriented electrical steel sheet |
| JP5332707B2 (en) * | 2009-02-20 | 2013-11-06 | 新日鐵住金株式会社 | Method for producing grain-oriented electrical steel sheet with extremely excellent magnetic properties |
| JP5684481B2 (en) * | 2010-02-15 | 2015-03-11 | 新日鐵住金株式会社 | Method for producing grain-oriented electrical steel sheet |
| WO2011102456A1 (en) * | 2010-02-18 | 2011-08-25 | 新日本製鐵株式会社 | Manufacturing method for grain-oriented electromagnetic steel sheet |
| US9761360B2 (en) * | 2012-03-29 | 2017-09-12 | Jfe Steel Corporation | Method of manufacturing grain oriented electrical steel sheet |
| CN103451515B (en) * | 2013-08-23 | 2015-09-23 | 安阳钢铁股份有限公司 | A kind of method controlling AlN inhibitor content in oriented silicon steel |
| KR101908045B1 (en) * | 2016-12-21 | 2018-10-15 | 주식회사 포스코 | Method for manufacturing grain oriented electrical steel sheet |
| JP6946846B2 (en) * | 2017-08-17 | 2021-10-13 | 日本製鉄株式会社 | Manufacturing method of grain-oriented electrical steel sheet |
| JP6946847B2 (en) * | 2017-08-17 | 2021-10-13 | 日本製鉄株式会社 | Manufacturing method of grain-oriented electrical steel sheet |
| CN116463545A (en) * | 2022-01-12 | 2023-07-21 | 宝山钢铁股份有限公司 | Copper-containing low-temperature oriented silicon steel and manufacturing method thereof |
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| US4898626A (en) * | 1988-03-25 | 1990-02-06 | Armco Advanced Materials Corporation | Ultra-rapid heat treatment of grain oriented electrical steel |
| JPH09118920A (en) * | 1995-10-25 | 1997-05-06 | Nippon Steel Corp | Stable manufacture of grain-oriented magnetic steel sheet excellent in magnetic property |
| FR2761081B1 (en) * | 1997-03-21 | 1999-04-30 | Usinor | METHOD FOR MANUFACTURING AN ELECTRIC STEEL SHEET WITH ORIENTED GRAINS FOR THE MANUFACTURE, IN PARTICULAR OF MAGNETIC CIRCUITS OF TRANSFORMERS |
| JP4203238B2 (en) * | 2001-12-03 | 2008-12-24 | 新日本製鐵株式会社 | Manufacturing method of unidirectional electrical steel sheet |
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