JP7801583B2 - Ferritic stainless steel sheet and its manufacturing method - Google Patents

Description

The present invention relates to ferritic stainless steel sheets for components that require high-temperature strength and workability, such as those used in automotive exhaust system parts, and to a method for manufacturing such sheets.

Efforts to reduce the weight of automobile bodies are being made as a measure to improve automobile fuel efficiency in response to global environmental issues. This involves increasing the strength of the steel plates used in automobiles and reducing their thickness, or by replacing them with aluminum or resin materials to reduce weight. The need for weight reduction extends not only to automobile bodies but also to various other parts. One example is the exhaust system components that process engine exhaust gases. Because exhaust system components are exposed to high-temperature engine exhaust gases, they require high-temperature strength and oxidation resistance. Therefore, heat-resistant stainless steels such as SUH409L, SUS436L, AISI429, and AISI441 are used, and exhaust system components weigh 20 to 30 kg per vehicle. In particular, the exhaust manifolds directly connected to the engine are exposed to high-temperature engine exhaust gases at the upstream end. Previously, they were cast, but now they are instead welded from stainless steel pipes or pressed and then welded from plates, resulting in weight reduction.

Parts exposed to high-temperature environments, such as exhaust manifolds, are called hot-end (H/E) parts, and Nb, Mo, Cu, and other elements are added to increase their high-temperature strength. However, solid-solution strengthening by these elements reduces workability at room temperature, so many technological developments have been made to date.

Patent Document 1 reports a technology in which ferritic stainless steel containing 0.4 to 1 wt% Nb is finish-annealed in the temperature range of 1100 to 1200°C to reduce the grain size to GSN No. 6 or less, thereby lowering the yield stress at room temperature and improving workability, and also increasing the high-temperature yield strength at 900°C. However, workability at room temperature is more influenced by the Lankford value than by the yield stress, and the coarsening of grains caused by finish-annealing at high temperatures of 1100°C or higher poses the problem of reducing thermal fatigue life due to the formation of orange peel during processing.

Patent Document 2 also discloses a method for producing ferritic stainless steel sheet with excellent high-temperature strength and formability, in which ferritic stainless steel containing 0.05-0.5% Ti and 0.1-1.0% Nb is cold-rolled at a reduction rate of 50-70% using mill rolls with a diameter of 100 mm or more, without hot-rolling annealing. This production method is commonly used today when improving the formability of Nb-added steel sheet for automobile exhaust systems. However, this production method alone does not provide sufficient formability for the high-process applications required today.

Therefore, in order to improve workability through techniques other than manufacturing methods, various steel types have been developed that combine the addition of Cu with the conventional addition of Nb and Mo to improve high-temperature strength. Patent Document 3 discloses a ferritic stainless steel for gas turbine exhaust gas path components containing 0.15-0.80% Nb and 1.0-3.0% Cu by mass. While Cu does not have as much of an effect on high-temperature strength as Nb or Mo, in environments where heating and cooling are repeated, such as automotive exhaust system components, it extends thermal fatigue life by repeating solution and precipitation. However, Cu significantly impairs pickling properties during hot-rolled sheet pickling, and also poses several issues, such as increased alloy costs and reduced recyclability. Furthermore, by reducing the amount of Nb and Mo and substituting them with Cu, the yield stress at room temperature is slightly reduced, making the material easier to work with, but the Lankford value is not improved, and the workability limit is not improved.

Patent Document 4 presents a technology that once again tackles improving workability through process optimization. It states that by precipitating many precipitates before finish annealing and then solutionizing them during finish annealing, both high-temperature strength and room-temperature workability can be achieved. By precipitating 0.4 to 1.2 mass% of fine precipitates before finish annealing and reducing the number of precipitates with a grain size of 0.5 μm or less to 0.5 mass% or less after finish annealing, it is said that the strength of the (222) plane is increased and workability is improved. However, heat treating hot-rolled sheet at 450 to 750°C, as in this invention, is not only inefficient as it extends the manufacturing period, but also reduces the toughness of the hot-rolled sheet and increases the risk of sheet fracture during the cold rolling process, making it an unrealistic process.

Japanese Patent Application Publication No. 05-331551 Japanese Patent Application Publication No. 06-179921 Japanese Patent Application Laid-Open No. 2002-004011 Japanese Patent Application Laid-Open No. 2002-194507

As such, it is difficult to economically impart higher workability than ever before to ferritic stainless steel sheets for automotive exhaust system parts containing Nb, Mo, and Cu, which have excellent high-temperature strength, using technologies disclosed to date, and new technological developments are needed.

Therefore, the present invention aims to provide a ferritic stainless steel sheet that has excellent high-temperature strength and workability at room temperature without compromising the high-temperature strength and productivity of ferritic stainless steel containing 0.2 mass% or more Nb, and a method for manufacturing the same.

The gist of the present invention is as follows.
(1) In mass%, C: 0.010% or less, Si: 0.05 to 1.20%, Mn: 0.05 to 1.50%, P: 0.035% or less, S: 0.010% or less, Cr: 10.5 to 20.0%, Ni: 0.01 to 0.60%, Cu: 0.01 to 1.60%, Mo: 0.01 to 2.0%, Al: 0.001 to 0.10%, Nb: 0.20 to 0.60%, N: 0.002 to 0.02%, the balance being Fe and impurities, and in a cross section in the rolling direction and thickness direction of the steel sheet, the number density of precipitates having an average major axis and minor axis of 0.05 to 0.2 μm is 0.2 particles/μm ² or less, and a grain size of 6 to 8 in GSN.
(2) The ferritic stainless steel sheet according to claim 1, further comprising, in mass %, one or more of the following elements from groups A to D: Sn: 0.001 to 0.2%, Co: 0.001 to 0.1% as A group elements, Ti: 0.003 to 0.15%, V: 0.005 to 0.10%, Zr: 0.005 to 0.10% as B group elements, B: 0.0003 to 0.0030% as C group elements, and Ca: 0.0001 to 0.0010%, REM: 0.001 to 0.020% as D group elements, in place of a portion of the Fe.

(3) A method for producing a ferritic stainless steel sheet according to (1) or (2), characterized in that a slab having the chemical composition according to (1) or (2) is cast to a thickness of 150 to 250 mm by continuous casting, heated to 1050 to 1250°C in a hot rolling heating furnace, hot rolled to a thickness of 3 to 8 mm, coiled at 350 to 650°C, pickled without or after hot-rolled sheet annealing, and subsequently cold-rolled to a thickness of 0.8 to 2.5 mm, and then annealed by heating to a temperature range of 840 to 1000°C at a heating rate of 50°C/s or more and holding at 1000 to 1060°C for 4 to 60 seconds.

The present invention provides a ferritic stainless steel sheet and its manufacturing method that has excellent high-temperature strength and workability at room temperature without compromising the high-temperature strength and productivity of ferritic stainless steel containing 0.2 mass% or more Nb.

By appropriately controlling the heating conditions during the finish annealing of ferritic stainless steel sheet containing 0.2% or more Nb by mass, the precipitation of fine precipitates of 0.05 μm or more and 0.2 μm or less is suppressed and the development of {111} orientation is promoted. Furthermore, to effectively utilize existing equipment and minimize new capital investment while suppressing the precipitation of harmful fine precipitates, the temperature is raised at a heating rate of 50°C/s or more between 840°C and 1000°C, the temperature range in which fine precipitates precipitate, during cold-rolled sheet annealing. Thereafter, annealing is performed for 2 to 60 seconds in the temperature range of 1000°C to 1050°C to achieve a uniform recrystallized structure with a grain size of 6 to 8 on the GSN scale. This makes it possible to impart higher workability at room temperature than ever before to ferritic stainless steel sheet containing 0.2% or more Nb. Not only can it be formed into thinner plates to reduce the weight of automobiles, but it can also be used to process parts with complex shapes optimized for improved fuel efficiency without compromising high-temperature strength. Furthermore, because there is no need to add large amounts of Cu or Mo, it is possible to reduce raw material costs and pickling costs, resulting in significant economic benefits.

This is a diagram showing the relationship between the heating rate and the Lankford value when a cold-rolled sheet of AISI429 (0.4% Nb) steel having a thickness of 1.5 mm is heated from room temperature to 1040°C at different heating rates and then annealed at 1040°C for 6 seconds. 2A, 2B, and 2C are photographs of precipitates observed by SEM in the cold-rolled annealed sheet shown in FIG. 1 for which the Lankford value was measured, where (A), (B), and (C) are photographs for the cases where the heating rates were 100, 20, and 3° C./s, respectively. FIG. 1 is a graph showing the relationship between the rapid heating start temperature and the number density of precipitates. FIG. 1 is a graph showing the relationship between the number density of precipitates and the average Lankford value. These are photographs of precipitates observed by SEM in a 1.2 mm thick cold-rolled sheet of AISI429 (0.4% Nb) steel, which was heated from room temperature to a predetermined rapid heating start temperature at 3°C/s, then rapidly heated to 1040°C at 100°C/s, and annealed at 1040°C for 6 seconds. (A) and (B) show the cold-rolled and annealed sheets when the rapid heating start temperatures were 950°C and 840°C, respectively, and (C) shows the cold-rolled sheet before annealing. FIG. 6 is a diagram showing the relationship between the rapid heating start temperature and the Lankford value when the rapid heating start temperature is changed to four different values by performing the same process as that shown in FIG. 5 .

The present invention will be described in detail below.
As described above, the present invention provides a ferritic stainless steel sheet with excellent high-temperature strength and workability, which is imparted with high workability at room temperature without impairing the high-temperature strength and productivity of ferritic stainless steel containing 0.2 mass% or more of Nb, and a method for producing the same. As an index of workability at room temperature, the target total elongation in a tensile test at room temperature is 30% or more. Regarding high-temperature strength, the target tensile strength at 700°C is 165 MPa or more.

The inventors observed precipitates in ferritic stainless steel AISI429 (0.4% Nb by mass) for automobile exhaust manifolds at 5000x magnification using an SEM. They observed numerous fine precipitates of 0.2 μm or less and a small number of relatively coarse precipitates of approximately 0.5 to 2.0 μm, with the fine precipitates forming rows. Given that these fine precipitates were not observed in the cold-rolled sheet before cold-rolling and annealing, but were observed after annealing, and given their row-shaped precipitation morphology, they are believed to have precipitated during annealing heating at the ferrite grain boundaries expanded by rolling.

Hereinafter, for the Lankford value, tests were performed in three directions: the L direction (r ₀ ) parallel to the rolling direction, the C direction (r ₉₀ ) parallel to the plate width direction, and the D direction (r ₄₅ ) between them, and the average Lankford value r _m was calculated as r _m = (r ₀ +r ₉₀ +2r ₄₅ )/4.

Therefore, cold-rolled sheets with the above composition were heated from room temperature to 1,040°C at heating rates of 3, 20, and 100°C/s, and annealed for 6 seconds. The Lankford values of the cold-rolled and annealed sheets were evaluated and shown in Figure 1. As a result, as shown in Figure 1, the material annealed at a heating rate of 100°C/s had a higher Lankford value than the other annealed materials.

Figures 2(A), (B), and (C) are SEM photographs of precipitates in a cross section (L cross section) of a cold-rolled annealed steel sheet in the rolling direction and thickness direction, respectively, when the heating rates were 100, 20, and 3°C/s. The horizontal direction of the photograph corresponds to the rolling direction, and the vertical direction corresponds to the thickness direction. The number density of fine precipitates with average major and minor diameters of 0.2 μm or less was evaluated in these SEM photographs. As shown in Figure 2(A), it was found that there were almost no fine precipitates at a heating rate of 100°C/s. In other words, the Lankford value was increased by suppressing the precipitation of fine precipitates during the heating process through rapid heating. This is thought to be due to the increased development of {111}-oriented grains, known as γ-fiber, which are advantageous for improving workability. Here, {111}-oriented grains refer to crystal grains whose {111} plane forms an angle of 15° or less with the steel sheet surface.

General industrial continuous annealing furnaces for steel strips use radiant heating from gas burners or radiant tubes to heat up, and the heating rate is 0.5 to 20°C/s for thicknesses of 1.5 to 2.0 mm, depending on the thickness. Using induction heaters or electric heating methods allows for more rapid heating, but rapid heating across the entire temperature range results in greater temperature variation across the strip width, which can lead to uneven material quality. Furthermore, this increases equipment costs.

Therefore, we considered incorporating rapid heating equipment into an existing steel strip annealing furnace to rapidly heat only a specific temperature range. Heating was performed from room temperature to the specified rapid heating start temperature T1 at 3°C/s, and then at 100°C/s to 1040°C, where it was held for 4 seconds. Rapid heating start temperatures T1 of 840, 880, 920, and 950°C were used. The precipitate density and average Lankford value were evaluated for cold-rolled annealed sheets annealed under each condition. Precipitate density was evaluated by selectively dissolving the matrix using the SPEED method to expose the precipitates in the cross section (L cross section) of the steel sheet in the rolling direction and thickness direction, and the precipitates were then observed using an FE-SEM.

Figure 3 shows the rapid heating start temperature on the horizontal axis and the precipitate number density on the vertical axis, with the number density of fine precipitates with average major and minor diameters of 0.2 μm or less indicated by ○ and the number density of precipitates with average major and minor diameters of more than 0.2 μm indicated by △. As a result, rapid heating from 840°C significantly reduced the number of fine precipitates, as shown in Figure 3.

In Fig. 4, the horizontal axis represents the number density of fine precipitates with average major and minor axis values of 0.2 µm or less, and the vertical axis represents the average Lankford value r _m . As shown in Fig. 4, the lower the number density of fine precipitates, the higher the average Lankford value r _m .

Based on the above results, the temperature range for rapid heating was determined to be 840°C or higher. Similar tests were conducted on the end temperature of rapid heating, and it was found that rapid heating should be performed up to 1000°C. It was also found that if the average heating rate in the temperature range for rapid heating is 50°C/s or higher, the precipitation of harmful precipitates can be suppressed and the Lankford value can be improved.

The grain size of steel sheet can be measured using the GSN based on JIS G0551 "Steel - Grain Size Microscope Test Method." After rapid heating, a uniform recrystallized structure is achieved, and in order to avoid roughness on the processed surface and achieve a high Lankford value, the steel sheet's grain size must be between GSN 6 and 8. Coarse grains with a GSN of less than 6 raise concerns about orange peel during processing. On the other hand, if the GSN exceeds 8, the elongation at room temperature will be low and the Lankford value will also be low.

It was found that when the heating and holding temperature after cold rolling and annealing is less than 1000°C, some unrecrystallized grains remain, and when the temperature exceeds 1100°C for 60 seconds, the GSN becomes smaller than 6. To achieve an appropriate range for the grain size of the recrystallized ferrite grains, it is necessary to optimize the temperature and time, but by holding the temperature in the 1020-1060°C range for 4-60 seconds, it is possible to achieve a GSN of 6-8. While controlling grain size has been achieved with existing technology, this invention suppresses the precipitation of harmful fine precipitates through rapid heating, making it possible to obtain the highest level of Lankford value for the same GSN. This is thought to be due to the further development of {111} oriented grains.

<Ferritic stainless steel plate>
Based on the above findings, the present inventors have discovered the ideal precipitate form and amount for ferritic stainless steel in the above applications, as well as a method for controlling the precipitation.

[Chemical composition]
The reasons for limiting each component are explained below. In the following explanation, "%" indicating the content of each element means "% by mass" unless otherwise specified.

C: 0.010% or less Solute C forms Cr carbides, which causes sensitization and reduces corrosion resistance, and requires the addition of stabilizing elements Nb, Ti, V, and Zr in proportion to the amount of C, which increases alloy costs. Therefore, the content must be 0.010% or less, and preferably 0.008% or less.
On the other hand, C is incorporated into molten steel during the process of smelting iron ore. While most of it can be removed during converter refining and degassing processes, this increases the refining time, impairing productivity, and also secondary oxidation of Cr reduces alloy yield. Therefore, it is preferable to set the C content at 0.001% or more. Considering the balance of strength and ductility required for sheet products, it is desirable to set the C content at 0.002% or more and 0.006% or less.

Si: 0.05-1.20%
The Si content is set to 0.05% or more because it is necessary for deoxidation during melting and refining and is also effective in suppressing the formation of oxide scale during hot rolling heating. To ensure the oxidation resistance required for automobile exhaust system parts, the Si content is preferably set to 0.10% or more.
On the other hand, since Si reduces the ductility of the steel sheet due to solid solution strengthening, the content is set to 1.20% or less. Furthermore, since Si forms an Si oxide film during annealing of the cold-rolled sheet and reduces the pickling property, the content is preferably set to 1.10% or less.

Mn: 0.05-1.50%
Mn is an element added as a deoxidizer and also contributes to increasing high-temperature strength in the medium temperature range. Furthermore, Mn-based oxides are formed on the surface during long-term use, and Mn is an element that contributes to the adhesion of scale (oxides) and the suppression of abnormal oxidation. Therefore, the content is set to 0.05% or more.
On the other hand, excessive addition not only reduces the toughness of the hot-rolled sheet due to the precipitation of the γ phase (austenite phase), but also forms MnS, which reduces the corrosion resistance, so the upper limit is set to 1.50%.
In addition, taking into consideration high-temperature ductility, scale adhesion, and suppression of abnormal oxidation, 0.20 to 1.00% is preferable.

P: 0.035% or less P is an element contained as an impurity in alloys such as molten pig iron and ferrochromium, which are raw materials. Since it is an element that is harmful to hot workability and toughness, the content is set to 0.035% or less. Since P is also an element that reduces workability, it is desirable to set it to 0.030% or less. Furthermore, excessive reduction in P content requires the use of high-purity raw materials, which leads to increased costs, so it is preferable to set the lower limit of P to 0.010%.

S: 0.010% or less S is an element that has a small solid solubility in the austenite phase and segregates at grain boundaries, promoting a decrease in hot workability. Since the effect becomes significant when the content exceeds 0.010%, the content is set to 0.010% or less. The lower the S content, the fewer sulfide-based inclusions there are, improving corrosion resistance. However, reducing S content increases the desulfurization load and production costs, so the lower limit is preferably set to 0.001%. The preferred range is 0.001% to 0.008%.

Cr: 10.5-20.0%
Cr is an essential element for ensuring oxidation resistance and corrosion resistance. If it is less than 10.5%, these effects are not realized, and if it exceeds 20.0%, it causes a decrease in workability and a deterioration in toughness, so the content is set to 10.5 to 20.0%. Note that, considering the decrease in crevice corrosion resistance due to the structure of exhaust system parts and the decrease in high-temperature strength due to oxidation, 13.5 to 19.0% is preferable.

Ni: 0.01~0.60%
Ni is present as an inevitable impurity in the alloy raw materials of ferritic stainless steel, and is generally contained in the range of 0.01 to 0.20%. It is also an element effective in suppressing the progression of pitting corrosion, and this effect is stably exhibited when added at 0.01% or more, so the lower limit is set to 0.01%. On the other hand, adding a large amount may lead to material hardening due to solid solution strengthening, so the upper limit is set to 0.60%. Considering the alloy cost, a range of 0.05 to 0.40% is desirable.

Cu: 0.01-1.60%
Cu is often unavoidably contained due to contamination from scrap during melting, etc. However, if high-purity raw materials are used to reduce Cu content, it may promote active dissolution during pit growth, impairing corrosion resistance, so the lower limit is set to 0.01% or more. To further improve corrosion resistance, 0.03% or more is desirable. Cu is also sometimes added proactively to increase high-temperature strength, but excessive content reduces hot workability and corrosion resistance, so the content is set to 1.60% or less. Since Cu precipitation may cause a decrease in corrosion resistance, 1.20% or less is desirable.

Mo: 0.01~2.00%
Like Cu, Mo has the effect of suppressing active dissolution and inhibiting the progression of pitting corrosion, making it an effective element for improving corrosion resistance. To achieve this effect, the content is set to 0.01% or more. To obtain even higher corrosion resistance, 0.03% or more is preferable. On the other hand, excessive Mo increases strength through solid solution strengthening and impairs press formability, so the upper limit is set to 2.00% or less. To achieve both corrosion resistance and workability, 1.50% or less is preferable.

Al: 0.001-0.10%
Al is an effective element for deoxidation, and this effect is manifested at 0.001% or more, so the lower limit is set to 0.001%. In order to obtain the deoxidation effect in combination with Si and Mn, it is preferable to set the content to 0.005% or more.
On the other hand, Al increases the basicity of the slag, precipitates water-soluble inclusions (CaS) in the steel, and may reduce corrosion resistance, so the upper limit is set at 0.10%. Also, considering the reduction in elongation due to alumina-based non-metallic inclusions, it is preferable to set the content at 0.01% or less.

Nb: 0.20-0.60%
Nb is an element that improves high-temperature strength and thermal fatigue properties, and also improves corrosion resistance by suppressing the formation of a Cr-deficient layer. To obtain the high-temperature strength required for exhaust system components, 0.20% or more must be added. From the viewpoints of wear resistance and corrosion resistance, 0.30% or more is preferable. On the other hand, excessive addition is undesirable because it generates Fe ₃ Nb ₃ C and Laves phase (Fe ₂ Nb) during use as an exhaust system component, impairing the solid-solution strengthening ability of Nb at high temperatures. Therefore, the content is set to 0.60% or less. To achieve both high-temperature strength, corrosion resistance, and workability, it is preferable to set the content to 0.45% or less.

N: 0.002-0.02%
N is an element that forms Cr nitrides at grain boundaries and causes sensitization, resulting in a decrease in corrosion resistance. While N can be reduced by degassing in a vacuum, prolonged degassing is difficult due to the drop in molten steel temperature. Since there is an industrial limit to how much N can be reduced, the lower limit for N is set to 0.002% or more. On the other hand, adding a large amount of nitrogen reduces workability and corrosion resistance, and requires the addition of a large amount of a stabilizing element, so the upper limit for N is set to 0.02%. From the viewpoint of achieving both workability and corrosion resistance, it is preferable that the N content be 0.008% or more and 0.015% or less.

The ferritic stainless steel sheet of the present invention contains the above components, with the remainder being Fe and impurities. Furthermore, in addition to the above elements, the present invention may contain, in mass percent, one or more of the following group A elements in place of a portion of the Fe: Sn: 0.001-0.2% and Co: 0.001-0.1%; B group elements: Ti: 0.005-0.15%, V: 0.005-0.10%, Zr: 0.005-0.10%; C group elements: B: 0.0003-0.0030%; and D group elements: Ca: 0.0001-0.0010% and REM: 0.001-0.020%. The upper limit may be set using high-purity raw materials.

Sn and Co, which are group A elements, are elements that contribute to improving high-temperature strength and corrosion resistance.
Sn: 0.001-0.20%
Like Mo and Cu, Sn is an element that enhances corrosion resistance by suppressing the progression of pitting corrosion, so it is desirable to add it as needed. To achieve this effect, 0.001% or more is preferable. However, it is known that Sn concentrates under the oxide scale and causes hot rolling cracks and defects. Furthermore, long-term aging at 400 to 700°C can reduce the toughness of steel, and it is therefore desirable to reduce the amount of Sn as much as possible. Therefore, adding 0.20% or less is preferable.

Co:0.001~0.10%
Co, like Nb, Mo, and Cu, has the effect of increasing high-temperature strength, but since it is a relatively expensive element, it is desirable to add it as needed. To achieve this effect, 0.001% or more is preferable. However, since Co has the effect of promoting the precipitation of Nb precipitates, it is preferable to add 0.10% or less.

The B group elements Ti, V, and Zr are stabilizing elements that contribute to improving additional corrosion resistance and workability.
Ti: 0.005-0.15%
Ti, together with Nb, is an element that combines with C, N, and S to improve corrosion resistance, intergranular corrosion resistance, room-temperature ductility, and deep drawability. Therefore, the Ti content is determined based on the economically achievable reduction amounts of C, N, and S and the amount of Nb added. However, excessive Ti addition impairs room-temperature workability due to solid solution strengthening, so the upper limit is set to 0.15% or less. However, since Nb is a more expensive element than Ti, adding Ti alone as a stabilizing element is preferable, and therefore adding Ti as a supplement to Nb is preferred. Therefore, adding 0.005% or more of Nb is preferable. Furthermore, since Ti has a higher sulfide system performance than Nb and is an effective element for suppressing the occurrence of pitting corrosion, adding 0.01% or more of Ti is desirable.

V: 0.005-0.10%
V functions as a stabilizing element for forming carbonitrides, so it is desirable to add it as needed. To achieve this effect, it is preferable to set the content to 0.005% or more. On the other hand, adding a large amount of V may promote the formation of coarse carbides due to solidification segregation, which may reduce ductility and toughness, so it is preferable to set the content to 0.10% or less.

Zr: 0.005% to 0.10%
Zr functions as a stabilizing element for forming carbonitrides, so it is desirable to add it as needed. To achieve this effect, the content is preferably 0.005% or more. On the other hand, adding a large amount of Zr may promote the formation of coarse carbides due to solidification segregation, which may reduce ductility and toughness, so the content is preferably 0.10% or less.

B, a C group element, is an element that contributes to improving secondary workability.
B: 0.0003% to 0.0030%
B has the effect of increasing grain boundary strength through grain boundary segregation and improving secondary workability, so it may be added as needed, and in order to exert this effect, the lower limit is preferably set to 0.0003% or more. However, excessive addition impairs toughness and corrosion resistance due to the precipitation of _Cr2B and (Cr,Fe) ₂₃ (C,B) ₆ , so the upper limit is preferably set to 0.0030%.

D group elements Ca and REM are elements that contribute to refinement of the solidification structure by oxide metallurgy and improvement of corrosion resistance by reducing sulfides.
Ca: 0.0001-0.0010%
Ca is an element added for desulfurization, and is also mixed in through refractory corrosion and slag entrainment. To prevent defects and deterioration of corrosion resistance caused by sulfides, the Ca content is preferably 0.0001% or more. However, excessive addition increases the amount of Ca-containing oxides, which can lead to nozzle clogging and defects, so the upper limit is preferably set to 0.0010%.

REM: 0.001-0.020%
REM is an element added for desulfurization and to fix P in steel to prevent temper embrittlement. It may also be added for the purpose of refining the solidification structure and improving oxidation resistance. To achieve these effects in combination, the REM content is preferably 0.001% or more. However, excessive addition of REM leads to the formation of coarse oxides and sulfides, which can cause nozzle clogging and defects, so the upper limit is preferably set to 0.020%.

[The number density of precipitates having an average major and minor diameter of 0.05 μm or more and 0.2 μm or less in the L cross section of the stainless steel plate is 0.2 particles/ ^μm2 or less]
As shown in FIG. 4 above, in a cross section (L cross section) of a steel sheet in the rolling direction and thickness direction, the lower the number density of fine precipitates with average major and minor axis values of 0.2 μm or less on the horizontal axis, the higher the average Lankford value. If the number density of fine precipitates is 0.2 particles/ ^μm2 or less, the average Lankford value is sufficiently high. Based on this result, in the present invention, the average major and minor axis values in the L cross section of a stainless steel sheet are specified to be 0.05 μm or more, and the number density of precipitates with average major and minor axis values of 0.2 μm or less is specified to be 0.2 particles/ ^μm2 or less. The reason for specifying the lower limit of the precipitate size as 0.05 μm or more is that it is difficult to accurately reveal precipitates with a size of 0.05 μm or less when the precipitates are revealed by a method such as the SPEED method, which dissolves only the matrix without dissolving the precipitates, and then observed using an FE-SEM or the like.

In order to develop the {111} oriented grains necessary to improve the Lankford value during the final annealing process of cold-rolled sheet, it is necessary to promote the generation and growth of {111} oriented recrystallized grains. Nb (C, N), which does not dissolve even during hot-rolling heating and has average major and minor axis sizes of approximately 0.5 to 2.0 μm, is uniformly dispersed regardless of the steel sheet structure and does not affect crystal orientation control during final annealing.

On the other hand, precipitates with average major and minor axis diameters of 0.05 to 0.2 μm precipitate during the temperature rise process of finish annealing, such as Fe ₃ Nb ₃ C and Fe ₂ Nb, mainly precipitate at ferrite grain boundaries, inhibiting the formation and growth of {111}-oriented grains that nucleate at the grain boundaries. When the SPEED method was used to selectively dissolve the matrix to reveal the precipitates and the number density was measured by observing them with an FE-SEM, the Lankford value was found to be high if the number density of fine precipitates was 0.2 particles/μm ² or less. To accurately measure the number density, it is desirable to photograph multiple fields of view at a magnification of 5,000 to 10,000 times and to use an examination area of 2,000 μm ² or more.

<Manufacturing method>
A steel sheet having a composition within the range of the present invention is cast by a continuous casting method under known conditions into a slab having a thickness of 250 to 150 mm, and then hot-rolled and cold-rolled by known methods. Preferably, the slab is maintained at or heated to a temperature not lower than 150°C, the surface of the slab is ground with a grinder as necessary, and then heated to 1250 to 1050°C in a hot-rolling heating furnace, and hot-rolled into a hot-rolled steel strip having a thickness of 3 to 8 mm. The hot-rolled steel sheet is then coiled at 350 to 630°C, with or without annealing, and then pickled and cold-rolled to a thickness of 0.8 to 2.5 mm.

During finish annealing after cold rolling, the steel is rapidly heated in the temperature range of 840°C to 1000°C at an average heating rate of 50°C/s or more, annealed at 1000 to 1060°C for 4 to 60 seconds, and then subjected to temper rolling and pickling to produce a cold-rolled annealed steel sheet with a thickness of 0.8 to 2.5 mm. This produces a ferritic stainless steel sheet with a number density of 0.05 to 0.2 μm precipitates of 0.2 particles/ ^μm² or less. Temper rolling may be performed before pickling. The steel sheet produced by the present invention may be welded and formed into a stainless steel pipe for use in exhaust system components.

The manufacturing method will be described in detail below.
[After slab casting, insert into hot rolling furnace]
The casting thickness of the slab must be 150 mm or more to obtain the hot rolling reduction necessary to develop the structure through hot rolling. Considering productivity during continuous casting, 200 mm or more is preferable. On the other hand, if the slab is too thick, the solidification structure becomes coarse, particularly at 1/4 of the plate thickness to the center, and no matter how much the hot rolling reduction is increased, it becomes impossible to develop the hot-rolled sheet structure necessary for improved workability. Therefore, the thickness is set to 250 mm or less. To reduce product ridging, a thickness of 240 mm or less is preferable. It is preferable to perform surface treatment on the cast slab using a grinder or the like, if necessary, depending on the surface properties of the slab and the intended use of the final product. Furthermore, unlike carbon steel, the target steel does not undergo phase transformation, so it retains a coarse solidification structure. This poses a risk of cracking due to reduced toughness caused by coarse grains. Therefore, after casting, it is preferable to maintain the temperature at 150°C or higher until it is inserted into the hot rolling furnace.

[Slab heating]
Slabs are heated for hot rolling. This reduces deformation resistance and allows hot rolling from the original slab thickness to a hot-rolled steel strip with a thickness of 3 to 8 mm. If the heating temperature is too low, the thin oxide scale of stainless steel makes it prone to seizure marks on the work rolls during hot rolling. Furthermore, recrystallization between rough hot rolling and finish hot rolling develops the texture necessary for improving the workability of the product after finish hot rolling. However, if the heating temperature is too low, recrystallization during this period may be insufficient, resulting in reduced workability. Considering these effects, a heating temperature of 1050°C or higher is necessary. Since an increase in the Nb content in the steel delays recrystallization, a temperature of 1150°C or higher is preferable. On the other hand, if the heating temperature is too high, the strength of the slab decreases, causing it to sag in the hot rolling furnace, resulting in poor transport, or the skid pushing into the slab may cause scratches; therefore, a temperature of 1250°C or lower is necessary. Higher heating temperatures may cause coarsening of crystal grains recrystallized between rough hot rolling and finish hot rolling, hindering the desirable development of the hot-rolled sheet texture, so the heating temperature is preferably 1200°C or lower.

[Hot rolling and coiling]
In the hot rolling, the steel sheet is rolled to a thickness of 25 to 40 mm in rough hot rolling and then rolled to a thickness of 3 to 8 mm in finish hot rolling. The coiling temperature is preferably 350 to 630°C in order to avoid uneven precipitate formation in the longitudinal direction of the hot rolled coil, prevent scale growth, and increase the efficiency of pickling.

[Cold rolling]
Hot-rolled coils are mechanically descaled using shot blasting, bending rolls, grinding brushes, etc., and then descaled by immersion in acid. Pickling is preferably performed using a bath containing sulfuric acid as the main component at 50 to 90°C in terms of the resulting surface properties. Furthermore, in applications where ridging of the product is a problem, it is preferable to anneal the hot-rolled sheet at 850 to 1100°C before pickling. At temperatures below 850°C, recrystallization is insufficient, and the ridging reduction effect is not achieved. On the other hand, at temperatures above 1100°C, the scale grows thick, increasing the load on the pickling process. The pickled hot-rolled coil is then cold-rolled to a thickness of 0.8 to 2.5 mm. To develop a texture that enhances the workability of the final product, it is preferable to use a cold-rolling work roll diameter of 400 mm or more. Furthermore, the thickness of the cold-rolled sheet is adjusted to the thickness required for each application of the automobile exhaust system part. In order to obtain a desirable rolling texture after cold rolling, the thickness is preferably 2.5 mm or less, and in order to ensure the high-temperature strength required for exhaust system parts, the thickness is preferably 1.0 mm or more.

[Finishing annealing]
During the finish annealing of the cold-rolled sheet, the temperature is increased in the temperature range of 840 to 1000° C. at an average temperature increase rate of 50° C./s or more.

As described above, by rapidly heating in the temperature range of 840 to 1000°C at an average heating rate of 50°C/s or more, it is possible to control the number density of precipitates having average major and minor axis values of 0.05 μm or more and 0.2 μm or less to 0.2 particles/ ^μm2 or less. Harmful fine precipitates precipitate even at temperatures lower than 840°C and above 1000°C, but they are harmless because they hardly precipitate during the heating time required for heating in a typical continuous annealing furnace.

On the other hand, recrystallization will not be complete below 1000°C, and it is inefficient to perform annealing to uniform the grain size after rapid heating using a separate heating method. Therefore, it is desirable to rapidly heat the material in the temperature range of 840 to 1000°C at an average heating rate of 50°C/s or more, and then set the maximum heating temperature for finish annealing to 1020°C or higher.

If the average heating rate is less than 50° C./s, the precipitation of harmful fine carbides Fe ₂ Nb and Fe ₃ Nb ₃ C competes with the nucleation of recrystallized grains at the ferrite grain boundaries, preventing the formation of a texture advantageous for workability.

Ferrite grain boundaries are likely to become precipitation sites because they are ferrite grain boundaries that have recrystallized after rough hot rolling or during annealing of the hot-rolled sheet. Subsequent finish hot rolling and cold rolling cause the ferrite grains to elongate, and dislocations accumulate at the grain boundaries due to sliding deformation caused by rolling, creating stable precipitation sites for precipitates and promoting precipitation. At ferrite grain boundaries that have not been rolled after recrystallization, these harmful fine precipitates do not form in the short time required for heating. An average heating rate of 100°C/s or more is desirable to reliably suppress fine precipitates. On the other hand, increasing the average heating rate too much can increase temperature variation across the width of the steel strip, potentially resulting in non-uniform material properties, so an average heating rate of 200°C/s or less is desirable.

In the temperature range from 840 to 1000°C, which corresponds to the precipitation nose of harmful fine precipitates, _{Fe2Nb and Fe3Nb3C} , harmful fine precipitates precipitate in a short time and inhibit the formation of {111} texture. Therefore, it was decided to continue rapid heating up to 1000°C.

Next, to achieve a grain size of 6 to 8 on the GSN scale, it is necessary to hold the material in the temperature range of 1000 to 1060°C for 4 to 60 seconds. If the holding temperature is below 1000°C, recrystallization will not be complete, resulting in a sheet material with high strength and low ductility. On the other hand, if the holding temperature exceeds 1060°C, the grains will become coarse and roughness will occur on the processed surface, which is undesirable. Furthermore, a holding time of 4 seconds or more is desirable to allow the grains to grow to a uniform grain size after recrystallization. On the other hand, holding for a long period of time can cause roughness on the processed surface due to coarsening of the grains, so a holding time of 60 seconds or less is desirable.

The steel plate is then cooled to room temperature, but since no phase transformation occurs and precipitates form slowly during the cooling process, rapid cooling is not necessary, so the average cooling rate is preferably 2°C/s or more and 50°C/s or less.

<Experiment>
A 200 mm thick slab having a composition of 13.3% Cr-1% Si-0.2% Mn-0.005% C-0.009% N-0.42% Nb (mass%) was cast, the slab temperature before insertion into the heating furnace was set to 300°C, and then heated to 1200°C. The slab was hot rolled to a thickness of 5 mm at a finish hot rolling temperature of 850°C, cooled with steam and water, coiled at 500°C, pickled, and then rolled into a 1.2 mm thick ferritic stainless steel cold-rolled steel sheet using a tandem cold rolling mill with a work roll diameter of 800 mm. The steel sheet was annealed to prepare a test material. In the first stage of cold-rolled sheet annealing, the temperature was increased from room temperature (23°C) to 840, 880, 920, and 950°C at an average heating rate of 3°C/s. After reaching each of the temperatures of 840, 880, 920, and 950°C (rapid heating start temperature T1), the temperature was subsequently increased to 1040°C at an average heating rate of 100°C/s. After holding at 1040°C for 6 seconds, the sheet was cooled to 200°C at 20°C/s and then allowed to cool naturally. The precipitates in the cross section (L cross section) of the annealed sheet in the rolling direction and thickness direction were revealed by SPEED etching, and photographs were taken at 5000 magnifications using an SEM. A Lankford test was performed in three directions using JIS No. 13B test pieces. Figure 5 shows SEM photographs of precipitates in the finish-annealed sheet and the cold-rolled sheet before annealing (C) when the rapid heating start temperature was 950°C (A) and 840°C (B). The evaluation conditions using the SEM were the same as those in Figure 2. Figure 6 shows the measurement results of the Lankford value for each rapid heating start temperature T1.

In the case of rapid heating from 840°C (Figure 5(B)), very few fine precipitates with average major and minor diameters of 0.2 μm or less are found, whereas in the case of rapid heating from 950°C (Figure 5(A)), fine precipitates are found to have formed in locations thought to be prior ferrite grain boundaries. Furthermore, in the cold-rolled sheet before final annealing (Figure 5(C)), just like the material rapidly heated from 840°C (Figure 5(B)), very few fine precipitates with average major and minor diameters of 0.2 μm or less are found, indicating that these fine precipitates precipitated during the temperature rise process of final annealing.

As shown in Figure 6, rapid heating from 840°C results in a higher Lankford value than rapid heating from 950°C.

The present invention will now be explained in more detail using examples.

A 250 mm thick slab with the steel composition shown in Table 1 was cast. The slab temperature before entering the heating furnace was 250°C, and after heating to 1200°C, it was hot-rolled to a thickness of 5 mm at a finish hot-rolling temperature of 850°C. It was then cooled with steam and water, coiled at 550°C, inserted into the furnace, held there for 1 hour, and air-cooled. Hot-rolled sheet annealing was omitted, and the slab was mechanically descaled by shot blasting, then pickled with sulfuric acid to remove scale. It was cold-rolled using a work roll with a roll diameter of 400 mm to produce a cold-rolled sheet with a thickness of 1.2 mm, which was annealed under the annealing conditions shown in Table 2.

The properties of the stainless steel sheets obtained in this manner were evaluated, and the results are shown in Table 2. In Tables 1 and 2, items that deviate from the scope of the present invention, items that deviate from the preferred manufacturing conditions of the present invention, and items that do not achieve the target quality of the present invention are underlined.

[Size and number density of precipitates]
The precipitates on the L cross section were revealed by SPEED etching, and after photographing them at 5000x magnification using an FE-SEM, the size and number density of the precipitates were measured. The average value of the minor axis and major axis was used as the size of the precipitates. Precipitates with a size of 0.05 to 0.2 μm and a number density of 0.2 particles/ ^μm2 or less were considered to be non-defective.

[Grain size]
The GSN was measured based on JIS G0551 "Steel - Grain size microscope test method." A GSN of 6 to 8 was considered a good product.

[Room temperature tensile test (elongation)]
Based on JIS Z2241 "Method of tensile testing for metallic materials," a tensile test was carried out at room temperature using a JIS No. 13B test piece in the L direction parallel to the rolling direction. Products with a total elongation of 30% or more in the tensile test were deemed to be good products.

[Lankford test]
Based on JIS Z 2254 "Testing method for plastic strain ratio of thin metal plate materials," a JIS Z 2201 No. 13B test piece was used to determine the plastic strain ratio (Lankford value) from the change in sheet width and the change in gauge length in the tensile direction after approximately 15% tension. Tests were conducted in three directions: the L direction ( _r0 ) parallel to the rolling direction, the C direction ( _r90 ) parallel to the sheet width direction, and the D direction ( _r45 ) between them. The average Lankford value _rm was calculated as _rm = ( _r0 + _r90 + _2r45 )/4.
Products with an average Lankford value _rm of 1.4 or more were considered to be good products.

[High temperature tensile test]
A high-temperature tensile test was carried out in the L direction at 700°C based on JIS G 0567 "High-temperature tensile test method for steel materials and heat-resistant alloys." Products with a tensile strength of 165 MPa or more at 700°C were considered to be good products.

In the method of the present invention, the average minor and major axis values observed on the L cross section of the annealed sheet after finish annealing were 0.2 μm or less, and the number density of fine precipitates of 0.05 μm or more was 0.2 particles/ ^μm2 or less. Compared to the comparative method, the elongation, Lankford value, and 0.2% yield strength at 700°C were higher. The GSN was also between 6 and 8, and it was confirmed that no orange peel occurred even after forming.

On the other hand, the comparative examples R1 to R7 had temperature histories during annealing that were outside the scope of the present invention, resulting in the precipitation of numerous fine precipitates, resulting in lower Lankford values than those of the present invention; or coarse grains with a GSN of less than 6 raised concerns about orange peel during processing; or a GSN of more than 8 resulted in low elongation at room temperature and a low Lankford value.

In R8, the N content was less than 0.002%, so that grain growth during the final annealing proceeded too much, resulting in a GSN of less than 6.0.
In R9, the C content exceeded 0.01%, so the amount of carbonitrides precipitated increased too much, resulting in low elongation at room temperature.
R10 had less than 0.05% Si and insufficient deoxidation during refining, resulting in an increase in oxide-based inclusions. In addition, the P content exceeded 0.035%, resulting in reduced elongation due to solid solution strengthening by P.
Since R11 contains more than 1.2% Si, the elongation at room temperature is reduced due to solid solution strengthening by Si.

In R12, the Mn content was less than 0.05%, so deoxidation was insufficient and many oxide-based inclusions were present, and furthermore, the N content exceeded 0.02%, so elongation decreased due to solid solution strengthening.
Since R13 contains more than 1.5% Mn, the elongation at room temperature is reduced due to solid solution strengthening by Mn.

In R14, the Cr content was less than 10.5%, and therefore the yield strength at 700°C was reduced.
Since R15 contains more than 20.0% Cr, the rolling texture during cold rolling is not sufficiently developed, resulting in a low Lankford value.

In R16, the Ni content exceeded 0.6%, which resulted in reduced elongation due to solid solution strengthening by Ni, and the Nb content was less than 0.2%, which resulted in reduced high-temperature strength.
R17 contained less than 0.01% Ni, less than 0.01% Mo, and less than 0.01% Cu, and therefore had a reduced high-temperature strength at 700°C.

Since R18 contains more than 1.60% Cu, fine Cu-rich clusters are formed during the cooling process after annealing the cold-rolled sheet, resulting in a decrease in elongation.
Since R19 contains more than 0.60% Nb, the elongation is reduced due to the precipitation of intermetallic compounds containing Nb and solid solution strengthening by Nb.

The present invention makes it possible to efficiently manufacture ferritic stainless steel sheets that have high high-temperature strength and excellent workability at room temperature. Therefore, the present invention contributes to reducing the weight and extending the life of automotive exhaust system parts, fuel system parts, structural members, and other components.

Claims (3)

The alloy contains, in mass%, C: 0.010% or less, Si: 0.05 to 1.20%, Mn: 0.05 to 1.50%, P: 0.035% or less, S: 0.010% or less, Cr: 10.5 to 20.0%, Ni: 0.01 to 0.60%, Cu: 0.01 to 1.60%, Mo: 0.01 to 2.0%, Al: 0.001 to 0.10%, Nb: 0.20 to 0.60%, N: 0.002 to 0.02%, with the balance being Fe and impurities;
In a cross section of the steel sheet in the rolling direction and the thickness direction, the number density of precipitates having an average major axis and minor axis of 0.05 to 0.2 μm is 0.2 particles/ ^μm2 or less,
A ferritic stainless steel plate characterized by having a crystal grain size of 6 to 8 in terms of GSN. Furthermore, as A group elements, one or two of Sn: 0.001 to 0.20% and Co: 0.001 to 0.10%;
As B group elements, one or more of Ti: 0.005 to 0.15%, V: 0.005 to 0.10%, and Zr: 0.005 to 0.10%;
As C group elements, B: 0.0003 to 0.0030%;
As D group elements, Ca: 0.0001 to 0.0010%, REM: 0.001 to 0.020%, one or two of them;
2. The ferritic stainless steel sheet according to claim 1, further comprising one or more of Groups A to D. A slab having the chemical composition according to claim 1 or claim 2,
It is cast to a thickness of 150 to 250 mm using the continuous casting method,
After heating to 1050-1250°C in a hot rolling furnace, the sheet is hot rolled to a thickness of 3-8 mm and coiled at 350-630°C.
Omit or perform hot-rolled sheet annealing, and pickle the hot-rolled sheet;
Subsequently, cold rolling is performed to a plate thickness of 0.8 to 2.5 mm,
The method for producing a ferritic stainless steel sheet according to claim 1 or 2, characterized in that, when the cold-rolled sheet is subsequently annealed, it is heated to a temperature range of 840 to 1000 ° C at a temperature increase rate of 50 ° C/s or more and held at 1000 to 1060 ° C for 4 to 60 seconds.