JP2020158815A - Stable austenitic stainless steel sheet - Google Patents

Abstract

To provide a stable austenitic stainless steel sheet that has a good polishability and workability.SOLUTION: A stable austenitic stainless steel sheet, in which the chemical composition is, in mass%, C: 0.15% or less, Si: 0.4 to 2.5%, Mn: 0.5 to 2.0%, Cr: 15.0 to 20.0%, Ni: 10.0 to 14.0%, N: 0.01 to 0.15%, Mo: 0 to 3.0%, Cu: 0 to 1.5%, Nb: 0 to 0.15%, V: 0 to 0.15%, Ti: 0 to 0.30%, B: 0 to 0.010%, balance: Fe and impurities, the relationship between Ni equivalent and Cr equivalent satisfies [Ni equivalent≥0.1×(Cr equivalent)2-3×(Cr equivalent)+30], the stacking fault energy SFE is 37 mJ/m2 and below, the average crystal grain size of the austenite phase is 8.0 μm or less, and the yield stress difference ΔYS is 100 MPa or less.SELECTED DRAWING: None

Description

The present invention relates to a stable austenitic stainless steel sheet.

In the housing of smartphones, precision equipment parts, etc., members with high surface gloss may be required. Stainless steel has excellent corrosion resistance and is an effective material for these members, but good polishability is required in order to stably obtain members with high surface gloss.

For example, in Patent Documents 1 to 4, improvement in polishability is studied. Patent Document 1 discloses a method for manufacturing a mirror-finished stainless steel sheet for a curved mirror having a wrapping-finished surface gloss and excellent image quality. Further, Patent Document 2 discloses an austenitic stainless steel for press forming, which has improved polishability for mirror finishing.

Further, Patent Document 3 discloses a method for producing a stainless steel strip and a steel plate having excellent polishability. Further, Patent Document 4 discloses a method for producing a steel strip having few surface minute defects in the production of austenitic stainless steel, martensitic stainless steel, or a two-phase stainless steel having a ferrite phase and an austenitic phase. Has been done.

Japanese Unexamined Patent Publication No. 3-169405 Japanese Unexamined Patent Publication No. 9-3605 Japanese Unexamined Patent Publication No. 62-253732 Japanese Unexamined Patent Publication No. 2000-273546

In addition, in precision parts used in the medical field, it is necessary to prevent the adhesion of debris and debris during polishing or processing. In addition, it is required not to adversely affect the operation of precision equipment. Therefore, the stainless steel sheet used for the above-mentioned parts and the like is also required to be non-magnetic.

By the way, among stainless steels, when a certain amount of a ferrite phase, a martensite phase, and a work-induced martensite phase generated in a semi-stable austenitic stainless steel is present, the stainless steel becomes magnetic.

On the other hand, the stable austenitic stainless steel sheet does not form a work-induced martensite phase without forming a second phase after processing and molding at at least room temperature or higher. Therefore, stable austenitic stainless steel has no magnetism, that is, it is non-magnetic. Therefore, stable austenitic stainless steel sheets are useful for the above-mentioned parts and the like.

Further, the above-mentioned parts and the like need to be processed into a product shape with high dimensional accuracy. Therefore, good workability is also required.

However, although the stainless steel sheets disclosed in Patent Documents 1 to 4 have good polishability, there is a possibility that sufficient workability cannot be obtained in consideration of processing into a complicated product shape. In addition, there are some that become magnetic. Therefore, the stainless steel sheets disclosed in Patent Documents 1 to 4 do not have sufficient properties of polishability, workability, and non-magnetism when used for the above-mentioned parts, and there is room for improvement.

An object of the present invention is to solve the above problems and to provide a stable austenitic stainless steel sheet having good polishability and workability and being non-magnetic.

The present invention has been made to solve the above problems, and the gist of the following stable austenitic stainless steel sheet is.

(1) The chemical composition is mass%
C: 0.15% or less,
Si: 0.4-2.5%,
Mn: 0.5-2.0%,
Cr: 11.0 to 20.0%,
Ni: 10.0 to 14.0%,
N: 0.01-0.15%,
Mo: 0-3.0%,
Cu: 0-1.5%,
Nb: 0 to 0.15%,
V: 0 to 0.15%,
Ti: 0 to 0.30%,
B: 0 to 0.010%,
Remaining: Fe and impurities,
The relationship between the Ni equivalent and the Cr equivalent defined by the following equations (i) and (ii) satisfies the following equation (iii).
The stacking defect energy SFE calculated by the following equation (iv) is 37 mJ / m ² or less.
The average crystal grain size of the austenite phase is 8.0 μm or less.
A stable austenitic stainless steel sheet in which the difference ΔYS in yield stress between the rolling direction and the plate width direction is 100 MPa or less.
Ni equivalent = Ni + 30 × C + 0.5 × Mn ・ ・ ・ (i)
Cr equivalent = Cr + Mo + 1.5 x Si ... (ii)
Ni equivalent ≧ 0.1 × (Cr equivalent) ^2-3 × (Cr equivalent) +30 ・ ・ ・ (iii)
SFE (mJ / m ² ) = 2.2 x Ni-1.1 x Cr-13 x Si-1.2 x Mn + 6 x Cu + 32 ... (iv)
However, each element symbol in the above formula represents the content (mass%) of each element contained in the steel, and if it is not contained, it is set to zero.

(2) The stable austenitic stainless steel sheet according to (1) above, wherein the sum of the X-ray random intensity ratios of the {110} <112> and {110} <001> directions at the center position of the plate thickness is 12.0 or less. ..

(3) The chemical composition is mass%.
Mo: 0.2-3.0%,
Cu: 0.1 to 1.5%,
Nb: 0.01-0.15%,
V: 0.01-0.15%,
Ti: 0.01 to 0.30%, and B: 0.0003 to 0.010%,
The stable austenitic stainless steel sheet according to (1) or (2) above, which contains one or more selected from the above.

According to the present invention, a stable austenitic stainless steel sheet having good polishability and workability and being non-magnetic can be obtained.

FIG. 1 is a diagram showing an ODF having a cross section of φ2 = 45 ° in which the crystal orientation is displayed.

The present inventors have conducted various studies in order to provide a stable austenitic stainless steel sheet that has good polishability and workability and is non-magnetic. As a result, the following findings (a) to (c) were obtained.

(A) In order to improve the polishability of the austenitic stainless steel sheet, it is effective to make the crystal grains of the austenitic phase of the steel sheet fine. Then, in order to make the crystal grains of the austenite phase finer, it is effective to reduce the stacking defect energy.

When the stacking defect energy of the austenite steel sheet is reduced, a large number of deformed twins are introduced in addition to the working strain during hot rolling or cold rolling, so that recrystallization nucleation occurs during rolling or subsequent annealing. The number of sites will increase and graining will be achieved.

In addition, since the austenite steel sheet is made finer, martensitic transformation occurs during rolling, a stable austenite phase is formed, and there is no possibility of magnetism. Therefore, it can be applied to precision equipment, especially equipment in which the use of magnetic materials is restricted, such as medical equipment.

(B) Further, in order to improve the workability, it is effective to control the difference in yield stress in the rolling direction and the rolling width direction within a certain range. For this purpose, it is desirable to suppress the formation of the texture from the {110} <112> orientation to the {110} <001> orientation, which increases the anisotropy of the yield stress.

(C) The texture is a structure formed in the rolling process. Therefore, in the cold rolling step, when the crystal grains are made into fine grains, the texture is developed, and finally the texture is maintained as it is. Therefore, in addition to making the crystal grains of the austenite phase finer, it is desirable to make the crystal grains finer in advance in the hot rolling step from the viewpoint of suppressing the formation of the texture.

The present invention has been made based on the above findings. Hereinafter, each requirement of the present invention will be described in detail.

1. 1. Chemical composition The reasons for limiting each element are as follows. In the following description, "%" for the content means "mass%".

C: 0.15% or less When the C content is more than 0.15%, coarse Cr carbides are precipitated at the grain boundaries during the final annealing, and the processability and corrosion resistance are deteriorated. Therefore, the C content is set to 0.15% or less. C has the effect of increasing the strength of the steel sheet at low cost. Therefore, the lower limit of the C content is not particularly set, but the C content is preferably 0.03% or more.

Si: 0.4-2.5%
Si is an element that has a great effect of lowering the stacking defect energy. If the Si content is less than 0.4%, this effect cannot be sufficiently obtained. Therefore, the Si content is set to 0.4% or more. However, if the Si content exceeds 2.5%, coarse oxides are formed and the processability is significantly reduced. Therefore, the Si content is set to 2.5% or less. The Si content is preferably 0.8% or more, and more preferably 1.2% or more. The Si content is preferably 2.0% or less, and more preferably 1.8% or less.

Mn: 0.5-2.0%
Mn is a strong austenite-forming element and has the effect of lowering stacking defect energy. Since these effects cannot be sufficiently obtained when the content is less than 0.5%, the Mn content is set to 0.5% or more. However, if the Mn content exceeds 2.0%, the workability deteriorates. Therefore, the Mn content is set to 2.0% or less. The Mn content is preferably 0.8% or more. The Mn content is preferably 1.5% or less, and more preferably 1.2% or less.

Cr: 11.0 to 20.0%,
Cr is a basic element of stainless steel and has the effect of forming a metal oxide layer on the surface of the steel material and enhancing corrosion resistance. Therefore, the Cr content is set to 11.0% or more. However, since Cr is a strong ferrite stabilizing element, if the Cr content is excessive, δ ferrite is generated. This δ ferrite deteriorates the hot workability of the material and also deteriorates the cold workability by remaining at room temperature. Therefore, the Cr content is set to 20.0% or less. The Cr content is preferably 13.0% or more. The Cr content is preferably 19.0% or less.

Ni: 10.0 to 14.0%
Ni is an austenite-forming element and is an important element for stably obtaining an austenite phase at room temperature. Therefore, the Ni content is set to 10.0% or more. However, Ni is an element that raises the energy of stacking defects, and twins are difficult to form during processing, which makes it difficult to refine the particles. Therefore, the Ni content is set to 14.0% or less. The Ni content is preferably 10.5% or more. The Ni content is preferably 13.5% or less.

N: 0.01 to 0.15%
Like C, N is a solid solution strengthening element and contributes to improving the strength of steel. Further, N is combined with Nb, Ti, and V and precipitated as a fine Nb compound during hot rolling and annealing, and has an effect of suppressing recrystallization and grain growth. Therefore, the N content is set to 0.01% or more. However, when the N content exceeds 0.15%, a large number of coarse nitrides are generated in the steel sheet manufacturing process, and these coarse nitrides serve as fracture starting points and significantly deteriorate hot workability. , Make manufacturing difficult. Therefore, the N content is set to 0.15% or less. The N content is preferably 0.03% or more. The N content is preferably 0.1% or less.

In addition to the above elements, the steel sheet according to the present invention may contain one or more elements selected from Mo, Cu, Nb, V, Ti, and B.

Mo: 0-3.0%
Mo has the effect of improving the corrosion resistance of the material. Therefore, it may be contained as needed. However, Mo is extremely expensive, and if the Mo content exceeds 3.0%, the cost will increase. In addition, an increase in Mo content increases Cr equivalents and destabilizes austenite. Therefore, the Mo content is preferably 3.0% or less, preferably 2.0% or less, and more preferably 1.0% or less. On the other hand, in order to obtain the above effect, the Mo content is preferably 0.2% or more.

Cu: 0-1.5%
Cu is an austenite-forming element and is an element capable of adjusting the stability of the austenite phase. Therefore, it may be contained as needed. However, if the Cu content exceeds 1.5%, segregation occurs at the grain boundaries during the production process, and this grain boundary segregation significantly deteriorates the processability during aging, making production difficult. Therefore, the Cu content is preferably 1.5% or less, preferably 1.2% or less. On the other hand, in order to obtain the above effect, the Cu content is preferably 0.1% or more.

Nb: 0 to 0.15%
V: 0 to 0.15%
Ti: 0 to 0.30%
B: 0 to 0.010%
Since Nb, V, Ti, and B all have the effect of suppressing recrystallization and aiming for fine granulation, they may be contained as necessary. However, if these elements are excessively contained, the development of {110} <112> and {110} <001> orientations is promoted, and the workability is deteriorated. Therefore, Nb: 0.15% or less, V: 0.15% or less, Ti: 0.30% or less, B: 0.010% or less. On the other hand, in order to obtain the above effects, it is preferable that Nb: 0.01% or more, V: 0.01% or more, Ti: 0.01% or more, and B: 0.0003% or more.

In the chemical composition of the steel sheet according to the present invention, the balance is Fe and impurities. Here, the "impurity" is a component mixed with raw materials such as ore and scrap, and various factors in the manufacturing process when steel is industrially manufactured, and is allowed as long as it does not adversely affect the present invention. Means something.

1-1. Relationship between Ni Equivalent and Cr Equivalent In order to provide good workability, the steel sheet according to the present invention needs to have a stable austenite phase without causing transformation into work-induced martensite or the like during processing. The stable austenite phase will be described later in 2-3.

Therefore, the steel sheet according to the present invention satisfies the following formula (iii) in the relationship between the Ni equivalent and the Cr equivalent, which will be described later. This is because if this relationship is not satisfied, it may have a two-phase or three-phase structure of austenite and ferrite and / or martensite. In addition, the austenite phase is unstable and martensitic transformation may occur when processed.

The Ni equivalent is the sum of the content of an element having the same effect as the content of Ni, which is an austenite-forming element, multiplied by a coefficient, and is defined by the following formula (i). Similarly, the Cr equivalent is the sum of the content of Cr, which is a ferrite-forming element, and an element having a similar effect multiplied by a coefficient, and is defined by the following equation (ii).
Ni equivalent = Ni + 30 × C + 0.5 × Mn ・ ・ ・ (i)
Cr equivalent = Cr + Mo + 1.5 x Si ... (ii)
Ni equivalent ≧ 0.1 × (Cr equivalent) ^2-3 × (Cr equivalent) +30 ・ ・ ・ (iii)
However, each element symbol in the above formulas (i) to (iii) represents the content (mass%) of each element contained in the steel, and if it is not contained, it is set to zero.

1-2. Laminated defect energy The steel sheet according to the present invention has a laminated defect energy (hereinafter, also referred to as “SFE”) calculated by the following equation (iv) of 37 mJ / m ² or less in order to provide good workability. To do.
SFE (mJ / m ² ) = 2.2 x Ni-1.1 x Cr-13 x Si-1.2 x Mn + 6 x Cu + 32 ... (iv)
However, each element symbol in the above formula represents the content (mass%) of each element contained in the steel, and if it is not contained, it is set to zero.

If the SFE is more than 37 mJ / m ² , twins are not sufficiently introduced during rolling deformation, and good workability cannot be obtained. Thus, SFE is a 37 mJ / ^{m 2} or less, it is preferable to be 30 mJ / ^{m 2} or less, and more preferably, 25 mJ / ^{m 2} or less. Although the lower limit of SFE is not particularly defined, it is considered that the component range in the present invention is -3 mJ / m ² or more.

2. 2. Metallographic structure 2-1. Average crystal grain size of the austenite phase In the steel sheet according to the present invention, the average crystal grain size of the austenite phase is 8.0 μm or less in order to provide good polishability. The average crystal grain size of the austenite phase is preferably 6.0 μm or less, and more preferably 5.0 μm or less. Here, the lower limit of the average crystal grain size of the austenite phase is not particularly determined, but is generally 1.0 μm or more.

The average crystal grain size of the austenite phase can be measured by the following method. Specifically, a sample having a shape of a plate thickness × 15 mm length × 10 mm width of a cross section parallel to the rolling direction (L cross section) at the center of the steel sheet is cut out and measured by EBSD. At this time, the measurement visual field is 100 μm × 100 μm including the outermost layer of the plate thickness × 15 mm L cross section, and the front and back surfaces are each 2 visual fields, for a total of 4 visual fields.

Of the regions determined to have an fcc structure in the above measurement, the region surrounded by grain boundaries excluding the annealed twin boundary (Σ3 grain boundary) with an orientation difference of 15 ° or more is regarded as one crystal grain, and is viewed in the observation field. The average area S per crystal grain is calculated from the number of crystal grains contained in the area. From the average area, the austenite average particle size D obtained by the following formula (1) was calculated. A version of TSL OIM Analysis 7 or later is used for the analysis of grain boundaries.
D = (2S / π) ^0.5 ... (1)

2-2. Random intensity ratio of {110} <112> and {110} <001> orientations The fiber texture from the {110} <112> orientation to the {110} <001> orientation is the characteristic orientation of the rolled austenite texture. Is. These orientation groups have large anisotropy, and are orientations that increase the yield stress in the rolling width direction. Therefore, from the viewpoint of workability, it is necessary to suppress the development of the texture of these orientation groups as much as possible.

The degree of tissue development can be evaluated by the sum of the X-ray random intensity ratios. Here, the X-ray random intensity ratio is a test material obtained by measuring the X-ray intensity of a standard sample having no accumulation in a specific direction and a test material by an X-ray diffraction method or the like under the same conditions. It is a numerical value obtained by dividing the X-ray intensity of the standard sample by the X-ray intensity of the standard sample.

When the sum of the X-ray random intensity ratios of the {110} <112> and {110} <001> directions at the center position of the plate thickness exceeds 12.0, the yield stress in the rolling direction and the width direction, which will be described later, is increased. The difference ΔYS becomes more than 100 MPa, and the workability is lowered. Therefore, the total sum of the X-ray random intensity ratios of the {110} <112> and {110} <001> orientations at the center position of the plate thickness is set to 12.0 or less. The total sum of the X-ray random intensity ratios of the {110} <112> and {110} <001> orientations at the center position of the plate thickness is preferably 11.0 or less, and more preferably 10.0 or less.

The lower limit of the total sum of the X-ray random intensity ratios is not particularly determined, but if the total sum of the X-ray random intensity ratios is less than 5.0, a phase other than the austenite phase such as a ferrite phase is developed. Suggests. Therefore, it is generally considered that the total sum of the X-ray random intensity ratios is 5.0 or more.

The total sum of the X-ray random intensity ratios can be measured by the following method. Specifically, the random intensity ratios of the {110} <112> and {110} <001> orientations are a plurality of {200}, {311}, and {220} pole diagrams measured by X-ray diffraction. It may be obtained from a crystal orientation distribution function (referred to as Origination Distribution Function, ODF) representing a three-dimensional texture, which is calculated by a series expansion method based on a pole diagram.

FIG. 1 shows an ODF having a cross section of φ2 = 45 ° in which the crystal orientation of the present invention is displayed. Strictly speaking, the {110} <112> azimuth refers to the azimuth expressed by φ1 = 55 ° and Φ = 90 °. However, since measurement errors may occur due to test piece processing and sample setting, the maximum value in the range of φ1 = 50 to 60 ° and Φ = 85 to 90 ° is represented as the intensity ratio in this direction.

Strictly speaking, the {110} <001> azimuth refers to the azimuth expressed by φ1 = 90 ° and Φ = 90 °. However, since measurement errors may occur due to test piece processing and sample setting, the maximum value in the range of φ1 = 85 to 90 ° and Φ = 85 to 90 ° is represented as the intensity ratio in this orientation. Here, as the orientation of the crystal, the orientation perpendicular to the plate surface is usually indicated by (hkl) or {hkl}, and the orientation parallel to the rolling direction is indicated by [uvw] or <uvw>. {Hkl} and <uvw> are generic names for equivalent planes, and (hkl) and [uvw] refer to individual crystal planes. That is, since the fcc structure is targeted in the present invention, for example, (111), (-111), (1-11), (11-1), (-1-11), (-11-1). , (1-1-1-1), (-1-1-) planes are equivalent and indistinguishable. In such a case, these orientations are collectively referred to as {111}.

Since ODF is also used to display the orientation of a crystal structure having low symmetry, it is generally expressed as φ1 = 0 to 360 °, Φ = 0 to 180 °, and φ2 = 0 to 360 °, and each of them is represented by φ2 = 0 to 360 °. The bearing is displayed in [hkl] (uvw). However, since the present invention targets an fcc crystal structure having high symmetry, Φ and φ2 are expressed in the range of 0 to 90 °. Further, the range of φ1 changes depending on whether or not symmetry due to deformation is taken into consideration when performing the calculation, but in the present invention, φ1 is expressed as φ1 = 0 to 90 ° in consideration of symmetry. That is, a method is selected in which the average value of the same direction at φ1 = 0 to 360 ° is expressed on the ODF of 0 to 90 °.

In this case, (hkl) [uvw] and {hkl} <uvw> are synonymous. Therefore, for example, the X-ray random intensity ratio of (110) [1-12] of the ODF in the φ2 = 45 ° cross section shown in FIG. 1 is synonymous with the X-ray random intensity ratio of the {110} <112> orientation. is there.

Further, the sample for X-ray diffraction is prepared as follows. Polish to a predetermined position in the plate thickness direction by mechanical polishing, chemical polishing, etc., and finish to a mirror surface by buffing. After that, the strain is removed by electrolytic polishing or chemical polishing, and at the same time, the thickness center position is adjusted to be the measurement surface. Here, since it is difficult to accurately set the measurement surface at a predetermined plate thickness position, a sample is prepared so that the measurement surface is within a range of 3% with respect to the plate thickness centering on the target position. do it.

The standard sample should be one in which the texture is completely randomized by compressing and sintering stable γ-based stainless steel powder, and mechanical polishing, chemical polishing, buffing, etc. are performed in the same manner as above. Although it is common, X-ray diffraction profile data of an ideal random sample obtained from calculation may be used.

When the measurement by X-ray diffraction is difficult, a statistically sufficient number of measurements may be performed by the EBSP (Electron Backscattering Pattern) method or the ECP (Electron Channeling Pattern) method.

2-3. Stable austenitic phase The steel sheet according to the present invention is a stable austenitic stainless steel in which a stable austenitic phase is formed. Here, the stable austenitic stainless steel refers to an austenitic stainless steel in which the austenitic phase is stably present by room temperature processing such as polishing or molding.

Specifically, it is an austenitic stainless steel in which the total fraction of the ferrite phase and the martensite phase is 1.0% or less when a tensile strain of 10% is applied assuming machining. If it is not a stable austenitic stainless steel, it may become magnetic when it is formed into a part shape by press forming or bending.

After processing at room temperature, it is desirable to reduce the formation of phases other than the austenite phase as much as possible. Therefore, the total fraction of the ferrite phase and the martensite phase when a tensile strain of 10% is applied is more preferably 0.5% or less, and more preferably 0.1% or less.

The total fraction of the ferrite phase and the martensite phase described above is a fraction of the entire phase and is a volume fraction. In addition, the total fraction does not clearly distinguish between the ferrite phase and the martensite phase. Both phases have a crystal structure b. c. c. It is a general term for iron crystal phases that have a structure and magnetism. The total fraction can be measured using a ferrite meter for a test piece to which a tensile strain of 10% is applied. As for the measurement conditions, the longitudinal center position is measured once on the front and back using three or more test pieces, and the average value is used.

A predetermined amount of tensile strain can be applied by, for example, a tensile test, and the tensile test is performed in accordance with JIS Z2241: 2011.

3. 3. Mechanical Properties In the steel sheet according to the present invention, the difference ΔYS in the yield stress between the rolling direction and the sheet width direction, which is also an index for evaluating workability, is defined. Specifically, in a steel sheet, if the difference ΔYS in the yield stress between the rolling direction and the plate width direction is more than 100 MPa, due to the difference in the amount of springback and the like, for example, the overhang height of the circular portion after press forming. Molding defects such as variations in the thickness occur. The difference ΔYS in the yield stress between the rolling direction and the plate width direction is preferably 100 MPa or less, preferably 80 MPa or less, and more preferably 50 MPa or less. The difference ΔYS in the yield stress between the rolling direction and the plate width direction is the absolute value of the difference between the yield stress in the rolling direction and the yield stress in the plate width direction.

The difference ΔYS in the yield stress between the rolling direction and the plate width direction may also be measured by the above-mentioned tensile test. The test piece for measuring the yield stress in the rolling direction is cut out so that the longitudinal direction of the sample is the rolling direction. Further, the test piece for measuring the yield stress in the plate width direction is cut out so that the longitudinal direction of the test piece is the plate width direction.

In the steel sheet according to the present invention, the polishability is evaluated by the roughness after polishing, and when the roughness Ra (average arithmetic roughness) after polishing is 0.050 μm or less, it is judged that the polishability is good. Roughness Ra is measured by the following procedure. Specifically, a sample having a length of 100 mm, a width of 150 mm, and a thickness of 0.2 mm is prepared and polished under the conditions of a surface pressure of 8.0 N / cm ² , abrasive grains # 400 alumina, a rotation speed of 300 rpm, and a polishing time of 10 seconds. I do. Then, the roughness after polishing is measured according to JIS B 0601: 2001.

4. Manufacturing Method Hereinafter, a preferable manufacturing method for the steel sheet according to the present invention will be described. The steel sheet according to the present invention can obtain the effect as long as it has the above-mentioned structure regardless of the manufacturing method. For example, the steel sheet can be stably obtained by the following manufacturing method.

4-1. Casting A steel having the chemical composition according to the present invention is melted and cast by a conventional method to obtain a steel piece to be subjected to hot rolling. This steel piece may be made by forging or rolling a steel ingot, but from the viewpoint of productivity, it is preferable to manufacture the steel piece by continuous casting. Further, it may be manufactured by using a thin slab caster or the like.

4-2. Hot rolling 4-2-1. Finished Rolling Start Temperature The obtained steel pieces are subjected to hot rolling to obtain a hot-rolled steel sheet. In the steel sheet according to the present invention, the start temperature of finish rolling in hot rolling (hereinafter, simply referred to as "finish rolling start temperature" or "F0T") is set to a temperature lower than the conventional method. Specifically, the finish rolling start temperature is preferably 1070 ° C. or lower. When the finish rolling start temperature exceeds 1070 ° C., the strain accumulation in hot rolling becomes insufficient, and the average crystal grain size of the hot-rolled steel sheet becomes coarse. As a result, the average crystal grain size of the austenite phase after cold rolling, which will be described later, also exceeds 8.0 μm.

Therefore, the finish rolling start temperature is preferably 1070 ° C. or lower, more preferably 1050 ° C. or lower, and further preferably 1020 ° C. or lower. The lower limit of the finish rolling start temperature is not particularly set, but it is considered that the temperature is usually 850 ° C. or higher because the temperature of 850 ° C. or higher causes an increase in deformation resistance and the load on the rolling mill becomes significant.

4-2-2. Hot rolling final stage rolling reduction In hot rolling, the steel sheet passes through the rolls of a plurality of hot rolling mills, and the final stage of hot rolling is performed multiple times in the entire hot rolling process. The final stage in the rolling path. In the steel sheet according to the present invention, the reduction rate of the final stage in hot rolling (hereinafter, simply referred to as "the final stage reduction rate of hot rolling") is preferably 15% or more. This is because if the reduction rate of the final stage of hot rolling is less than 15%, recrystallization occurs in the final stage and a coarse processed structure remains.

Therefore, the final stage reduction rate of hot rolling is preferably 15% or more, more preferably 18% or more, and further preferably 20% or more. On the other hand, the upper limit of the final stage reduction rate of hot rolling is not particularly determined, but if it exceeds 40%, the load of the hot rolling mill may be significantly increased. Therefore, the final stage reduction rate of hot rolling is preferably 40% or less, and more preferably 35% or less.

4-2-3. Hot-rolled sheet annealing The obtained hot-rolled steel sheet may be annealed by hot-rolled sheet for the purpose of softening. However, when the annealing temperature exceeds 1150 ° C., grain growth occurs remarkably. Therefore, the temperature reached during annealing of the hot-rolled steel sheet (also simply referred to as "hot-rolled sheet annealing temperature") is preferably 1150 ° C. or lower, more preferably 1100 ° C. or lower, and 1000 ° C. or lower. Is even more preferable.

4-3. Cold rolling and final annealing Subsequently, cold rolling and annealing are repeated once or a plurality of times on the hot-rolled steel sheet to produce the cold-rolled steel sheet. At that time, it is desirable that the final reduction rate of cold rolling, the temperature reached in the final annealing (hereinafter, simply referred to as “the temperature reached in the final annealing”), and the holding time at the temperature, which will be described later, be within the ranges described below. Conditions other than the final cold rolling reduction rate, the final annealing temperature, and the holding time may be in accordance with a conventional method.

4-3-1. Final cold rolling reduction rate The final reduction rate for cold rolling is preferably 30 to 80%. When the final reduction rate of cold rolling is less than 30%, the strain introduced during cold rolling is insufficient and the number of recrystallized nucleation sites decreases. As a result, the average crystal grain size of the austenite phase of the cold-rolled steel sheet exceeds 8.0 μm. Therefore, the final rolling reduction ratio for cold rolling is preferably 30% or more, more preferably 35% or more, and even more preferably 40% or more.

On the other hand, when the final reduction rate of cold rolling is more than 80%, the processed texture develops remarkably, and the texture in the {110} <001> orientation develops. As a result, the anisotropy of the yield stress becomes large, which may lead to deterioration of workability. Therefore, the final rolling reduction rate for cold rolling is preferably 80% or less, more preferably 75% or less, and even more preferably 70% or less.

4-3-2. Final annealing reaching temperature and holding time The final annealing reaching temperature is preferably in the range of 750 to 1050 ° C. If the final annealing temperature is less than 750 ° C., recrystallization is not completed and the processed grains spread in the rolling direction remain, so that the average crystal grain size of the austenite phase does not finally become 8.0 μm or less. Therefore, the final annealing temperature is preferably 750 ° C. or higher, more preferably 780 ° C. or higher, and even more preferably 800 ° C. or higher.

On the other hand, when the final annealing temperature exceeds 1050 ° C., grain growth occurs and the crystal grains become coarse. Therefore, the final annealing temperature is preferably 1050 ° C. or lower, more preferably 1000 ° C. or lower, and even more preferably 970 ° C. or lower.

Further, the holding time at the final annealing temperature is preferably in the range of 100 s or less. When the holding time at the final annealing temperature exceeds 100 s, grain growth occurs and the crystal grains become coarse. Therefore, the holding time at the final annealing temperature is preferably 60 s or less, more preferably 20 s or less, and further preferably 10 s or less.

Although the lower limit of the holding time is not particularly defined, it is preferable to set the lower limit to 0.1 s because it is difficult to control the holding time to 0.1 s or less.

After the final annealing, cold rolling for the purpose of adjusting the mechanical properties of the steel sheet, that is, temper rolling may be performed. After that, heat treatment may be performed for the purpose of reducing the residual stress that causes the shape change of the steel sheet without significantly changing the characteristics of the steel sheet, that is, for the purpose of removing strain.

It is desirable that the temper rolling ratio is 50% or less. This is because the temper rolling ratio can be adjusted to the required mechanical properties at 50% or less. Subsequently, the heat treatment temperature is preferably 600 ° C. to 900 ° C., more preferably 650 to 850 ° C. This is because the strain removing effect cannot be obtained at a temperature lower than 600 ° C. and reverse transformation does not occur. Further, at a temperature exceeding 900 ° C., the effect of performance adjustment in cold rolling disappears.

Hereinafter, the present invention will be described in more detail with reference to Examples, but the present invention is not limited to these Examples.

Steel pieces having the chemical composition shown in Table 1 are melted to produce steel pieces, and the steel pieces are heated to perform rough rolling hot, and then finish rolling under the conditions shown in Table 2 below. Was done. In Table 2, F0T [° C.] indicates the finishing hot spreading start temperature, and HA indicates the hot rolling plate annealing reaching temperature. After the finish rolling, the hot-rolled annealed sheet was pickled, then subjected to one intermediate cold rolling and intermediate annealing, and then subjected to final cold rolling.

In Table 2, CR [%] indicates the final rolling reduction ratio for cold rolling. After cold rolling, final annealing was performed. In the final annealing, as shown in Table 2, the temperature was raised to the temperature at which the final annealing was reached, and the temperature was maintained at that temperature for a predetermined time. Then, cooling was performed to obtain a steel plate. AT [° C.] indicates the final annealing temperature.

(Average crystal grain size of austenite phase)
The average crystal grain size of the austenite phase was measured for the obtained steel sheet. Specifically, a sample having a shape of thickness × 15 mm length × 10 mm width was cut out from a cross section (L cross section) parallel to the rolling direction at the center of the steel sheet, and measured by EBSD. At this time, the measurement visual field was set to 100 μm × 100 μm so as to include the outermost layer, and the front and back surfaces were each 2 visual fields, for a total of 4 visual fields.

Of the regions determined to have an fcc structure in the above measurement, a region surrounded by grain boundaries having an orientation difference of 15 ° or more and not twin grain boundaries (Σ3 grain boundaries) is regarded as one crystal grain and is viewed in the observation field. The average area S per crystal grain was calculated from the number of crystal grains contained in the area. From the average area, the austenite average particle size D obtained by the following formula (1) was calculated. TSL OIM Analysis 7 was used for the analysis of grain boundaries.
D = (2S / π) ^0.5 ... (1)

(Random intensity ratio of {110} <112> and {110} <001> orientations)
The random strength ratios of the {110} <112> and {110} <001> orientations of the steel sheet were measured as follows. First, the steel sheet was mechanically polished and buffed, then electropolished to remove strain, and X-ray diffraction was performed using a sample adjusted so that the center position of the sheet thickness was the measurement surface.

In addition, a standard sample prepared by a method such as sintering and having no accumulation in a specific orientation was prepared. The standard sample was prepared by mechanical polishing, chemical polishing, buffing, and the like. X-ray diffraction of the standard sample was also performed under the same conditions. Next, based on the {200}, {311}, and {220} pole figures obtained by X-ray diffraction, an ODF was obtained by a series expansion method. Then, the random intensity ratio was determined from this ODF. In Table 2, the random intensity ratios of the {110} <112> and {110} <001> orientations are simply referred to as the random intensity ratios.

(Total fraction of ferrite phase and martensite phase when 10% tensile strain is applied)
In order to evaluate whether it is a stable austenitic stainless steel, the total fraction of the ferrite phase and the martensite phase when a tensile strain of 10% was applied was measured by the following method. Specifically, a tensile test piece (JIS No. 13B) having a longitudinal direction in the rolling direction is prepared, a strain of 10% is applied by a tensile test in accordance with JIS Z2241: 2011, and then a ferrite meter is applied to the parallel portion (20 mm). Measured at. In Table 2, the total fraction of the ferrite phase and the martensite phase when a tensile strain of 10% is applied is simply referred to as the total fraction after 10% tensile.

(Difference in yield stress between rolling direction and plate width direction ΔYS)
The difference ΔYS in the yield stress between the rolling direction and the plate width direction was measured by performing a tensile test in accordance with JIS Z2241: 2011. The shape of the test piece was the above-mentioned shape. The test piece for measuring the yield stress in the rolling direction was cut out so that the longitudinal direction of the sample was the rolling direction. Further, the sample was cut out so that the longitudinal direction of the sample for measuring the yield stress in the plate width direction was the plate width direction. In Table 2, L-YS [MPa] indicates the yield stress in the rolling direction, and C-YS [MPa] indicates the yield stress in the plate width direction.

(Abrasiveness)
The polishability of the steel sheet was evaluated as follows. Specifically, a sample having a length of 100 mm, a width of 150 mm, and a thickness of 0.2 mm is prepared and polished under the conditions of a surface pressure of 8.0 N / cm ² , abrasive grains # 400 alumina, a rotation speed of 300 rpm, and a polishing time of 10 seconds. Was done. Then, the roughness after polishing was measured according to JIS B 0601: 2001. Then, when the roughness Ra (average arithmetic roughness) after polishing was 0.050 μm or less, it was judged that the polishability was good.

No. which is an example of the present invention. In 1, 2, 4, 6, 8, 10, 13, 15, and 18, the value of the average crystal grain size of the austenite phase was small, and the polishability was also good. Further, in the above example, the difference ΔYS in the yield stress between the rolling direction and the plate width direction was small, and the anisotropy of the yield stress was also small.

No. which is a comparative example. 19 to 23 are examples which do not satisfy the chemical composition defined in the present invention. No. In No. 19, the Si content was more than 2.5% and was excessive, so that the hot workability was significantly deteriorated and cracks were generated in the hot spreading stage, so that the test was stopped. No. In No. 20, the Si content was low, and in this case, the SFE was too high, so it is considered that twinning did not occur during deformation such as rolling. After that, the recrystallization nucleation site became insufficient, and the crystal grains of the austenite phase became coarse grains, resulting in inferior polishability.

No. In No. 21, the Cr content was below the specified range and the SFE was too high, so that the crystal grains of the austenite phase became coarse grains, resulting in poor polishability. No. In No. 22, recrystallization was remarkably suppressed, the texture was strongly developed, and the anisotropy of the yield stress was increased because the Nb content was excessive. No. In No. 23, the Ni content was below the specified range, the relationship between the Ni equivalent and the Cr equivalent was not satisfied, and a stable austenite phase was not formed. In other words, it did not become a stable austenitic stainless steel.

No. 3, 5, 7, 9, 11, 12, 14, 16, and 17 all satisfy the provisions of the present invention in terms of chemical composition, but sufficient polishability is not ensured, and / or yield stress. The isotropic property was not ensured, the difference in yield stress ΔYS was large, and the anisotropy was large.

No. In No. 3, since the final stage reduction rate of hot rolling was low, recrystallization did not occur at the stage of hot rolling, and a coarse rolled structure remained after that, and the crystal grains of the austenite phase became coarse grains. As a result, the polishability was inferior. No. In No. 5, since the temperature at which the hot-rolled plate was annealed was too high, grain growth occurred, and the crystal grains of the austenite phase became coarse grains. As a result, the polishability was inferior. No. In No. 7, since the final reduction rate of cold rolling was too low, recrystallization did not occur and the structure became a recovered structure, so that the crystal grains of the austenite phase became coarse grains. As a result, the polishability was inferior.

No. In No. 9, since the final reduction rate of cold rolling was too high, the textures in the {110} <112> and {110} <001> orientations developed, and the anisotropy of the yield stress became large. No. No. 11 is an example in which recrystallization did not occur and processed grains remained because the final annealing temperature was too low. In this example, the coarse grains expanded in the rolling direction remain, and the texture in the {110} <112> to {110} <001> orientations also develops, so that the abrasiveness deteriorates and the yield stress is anisotropy. Has also grown.

No. In No. 12, the final reduction rate of cold rolling is too high and the final annealing temperature is too low, so that recrystallization is completed and fine grains are achieved, but the texture is strong, so the anisotropy of the yield stress is large. became. No. In No. 14, since the final annealing temperature was too high, the crystal grains of the austenite phase became coarse grains, and the polishability was lowered. No. In No. 16, the austenite phase in the cold-rolled steel sheet was also coarse-grained because the hot-rolled sheet in which the accumulation of strain during hot-rolling was insufficient became coarse-grained because F0T was too high. As a result, the polishability was inferior. No. No. 17 became fine particles because F0T was too high and the final rolling reduction of cold rolling was too high, but the anisotropy of the yield stress became large.

As described above, the stable austenite steel sheet having fine particles satisfying the provisions of the present invention and having a small anisotropy of yield stress had good polishability and workability. On the other hand, in the comparative example which does not satisfy the provisions of the present invention, it is considered that transformation occurs when the coarse grain and / or the yield stress is highly anisotropy, the austenite phase is not stable, and strain is applied. There was an example. Therefore, the result is that at least one of polishability and processability is inferior.

Claims (3)

The chemical composition is mass%,
C: 0.15% or less,
Si: 0.4-2.5%,
Mn: 0.5-2.0%,
Cr: 11.0 to 20.0%,
Ni: 10.0 to 14.0%,
N: 0.01-0.15%,
Mo: 0-3.0%,
Cu: 0-1.5%,
Nb: 0 to 0.15%,
V: 0 to 0.15%,
Ti: 0 to 0.30%,
B: 0 to 0.010%,
Remaining: Fe and impurities,
The relationship between the Ni equivalent and the Cr equivalent defined by the following equations (i) and (ii) satisfies the following equation (iii).
The stacking defect energy SFE calculated by the following equation (iv) is 37 mJ / m ² or less.
The average crystal grain size of the austenite phase is 8.0 μm or less.
A stable austenitic stainless steel sheet in which the difference ΔYS in yield stress between the rolling direction and the plate width direction is 100 MPa or less.
Ni equivalent = Ni + 30 × C + 0.5 × Mn ・ ・ ・ (i)
Cr equivalent = Cr + Mo + 1.5 x Si ... (ii)
Ni equivalent ≧ 0.1 × (Cr equivalent) ^2-3 × (Cr equivalent) +30 ・ ・ ・ (iii)
SFE (mJ / m ² ) = 2.2 x Ni-1.1 x Cr-13 x Si-1.2 x Mn + 6 x Cu + 32 ... (iv)
However, each element symbol in the above formula represents the content (mass%) of each element contained in the steel, and if it is not contained, it is set to zero. The stable austenitic stainless steel sheet according to claim 1, wherein the sum of the X-ray random intensity ratios in the {110} <112> and {110} <001> directions at the center position of the plate thickness is 12.0 or less. When the chemical composition is mass%,
Mo: 0.2-3.0%,
Cu: 0.1 to 1.5%,
Nb: 0.01-0.15%,
V: 0.01-0.15%,
Ti: 0.01 to 0.30%, and B: 0.0003 to 0.010%,
The stable austenitic stainless steel sheet according to claim 1 or 2, which contains one or more selected from.

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