GB1574875A - Production methods for steel sheet containing carbide or nitride forming elements - Google Patents

Description

(54) PRODUCTION METHODS FOR STEEL SHEET CONTAINING CARBIDE OR NITRIDE FORMING ELEMENTS (71) We, NIPPON STEEL CORPORATION, a Japanese Company, of No. 6-3, 2-chome, Ote-machi, Chiyoda-ku, Tokyo, Japan, do hereby declare the invention, for which we pray that a patent may be granted to us, and the method by which it is to be performed, to be particularly described in and by the following statement:- This invention relates to a method for producing a low carbon steel strip or sheet containing carbide or nitride forming elements. The invention also relates to low carbon steel sheet or strip produced by such methods.

The usual production processes for steel strip or sheet utilise a starting material in the form of a steel slab, the slab being produced from an ingot by a process including blooming, or the slab being produced directly by a casting process, such as a continuous casting process. After manufacture, the steel slab usually is cooled down to the ambient temperature, but thereafter, the slab is heated up to a temperature in the range of 1200 to13000C for more than three hours in a slab reheating furnace. The hot slab is then fed to a hot rolling mill and hot rolled into the desired thickness.

Since the continuous casting technique has been developed, it has been recognised in the art that there would be advantages in hot rolling, a continuously cast steel slab directly after casting, i.e. without the necessity of reheating the slab. This process of directly rolling a cast steel slab (hereinafter called a "direct hot rolling" process) is well-known and well established and various ways of carrying out this process have been proposed. The main objects of a direct hot rolling process in the past has been to make the processing steps of casting and hot rolling continuous and to save energy as compared with a process in which the slab is cooled down to the ambient temperature and then reheated in a slab-reheating furnace prior to hot rolling.

There has been very little consideration of the technical problems to be solved in a direct hot rolling process and how the process would influence the quality of the final product, from a metallurgical viewpoint.

This invention stems from careful studies on the relationship between a heat diagram for a steel slab and that for a hot rolled steel strip, and an important relationship between the two steps has been established, particularly for steel containing carbide or nitride forming elements such as acid soluble aluminum (hereinafter called Sol Al), Ti, V and Nb.

According to this invention, there is provided a method for producing a rolled, low-carbon steel material, comprising forming a steel slab from a low-carbon steel composition containing (by weight) at least one carbide or nitride forming element selected from 0.015 to 0.10% acid-soluble aluminium, 0.01 to 0.10% titanium, 0.01 to 0.15% vanadium, and 0.01 to 0.10% niobium, by a forming process in which the finished slab is at a temperature above the Ar3 point of the steel and in which the or each said element is dissolved and dispersed evenly throughout the slab, maintaining the temperature of the thus-formed slab at a temperature above the Ar3 point of the lowcarbon steel from the time of the formation of the slab until the commencement of rolling of the slab, and directly carrying out hot rolling of the slab while the temperature of the slab is above the Ar3 point of the low carbon steel.

The hot rolled steel strip or sheet is preferably cold rolled and annealed to complete the production of a low alloy steel material.

The process of this invention basically comprises hot rolling a slab of low-carbon steel which includes at least one carbide or nitride forming element, such as soluble aluminium (Sol Al), titanium, vanadium or niobium, the temperature of the slab being maintained above the Ar3 point while the slab is moved from the casting or blooming step to the hot rolling operation. It has been found that when this condition is maintained, carbides or nitrides - which are precipitated after hot rolling - are uniformly precipitated and finely dispersed in the hot steel during succeeding steps, these precipitates act effectively during the succeeding steps to raise the quality of the final product.

According to the prior art usual sheet or strip production process, the cast or bloomed steel slab is cooled down to the ambient temperature before being hot rolled. In the cold slab, the carbides and nitrides are completely precipitated, the precipitates growing to form large grains during the cooling. Therefore, in the prior art process, reheating at a high temperature for several hours is necessary in order to redissolve these precipitates and keep them in the dissolved state before starting the hot rolling operation. Typically, the reheating is performed above 1200"C for more than 3 hours. However, even if these precipitates are completely re-dissolved by this reheating step, each element is not completely and uniformly re-dispersed in the steel as in the dissolved state. Therefore, when the carbides or nitrides again precipitate in the succeeding steps, these precipitates do not re-precipitate uniformly, and are not so effective at yielding the desired qualities of the steel.

In a hot slab which has been produced by an ingot-making process or by a continuous casting process, each element is in a dissolved and uniformly dispersed state, and the present invention effectively utilizes this steel as a starting material for producing a low-carbon final product having good deepdrawing qualities or high strength.

The carbide or nitride precipitates having an important influence on the quality of steel are AIN, TiC, V(CN) and Nb(CN). The behaviour of these rrecipitates with respect to ingot casting, continuous castling, heating, hot-rolling and annealing, respectively, has been studied and from the results, the most favourable conditions for producing steel strips or sheets containing at least one of the elements selected from Sol Al, Ti, V and Nb have been determined.

In the method of this invention, carbide or nitride forming elements are kept in the dissolved state in the steel by keeping the steel at a temperature above the Ar3 point from the casting or blooming step to the start of the hot rolling, and then the slab is directly hot rolled without the temperature thereof being allowed to fall. If necessary, heat can be added to the slab so that a uniform temperature of the entire slab is maintained. Sueh heating can be carried out by heating the slab at a temperature below 12800C, and preferably -no higher than 1250"C, rather than by heating at a temperature much higher than the Ar3 point, as in a conventional high temperature slab reheating process.

The precipitates have many purposes.

One of these purposes is controlling the recrystallization texture, and another is controlling the recrystallized grain size and shape and, in addition, ensuring the steel material attains a certain strength. The most important aspects of these precipitates are the time at which they are formed, the physical form of the precipitates and their dispersion in the steel. For example, when a Al-killed steel is used for producing a cold rolled steel sheet having a good deep drawing quality, it is important to control the texture of the recrystallized grains so as to develop a recrystallized texture which is favourable for deep drawability. To achieve this, in a known process after hot rolling, Al and N are kept dissolved by, for instance, coiling the hot rolled strip at a low temperature, e.g. from 500 to 6500 C, and the Al and N are precipitated as AIN at the time of annealing after cold-rolling.

As a result of an extensive study of the conditions for precipitating AIN in Al-killed steel, a specific manner of achieving the most favourable conditions has been found, and this comprises feeding a slab to a hot rolling mill at a temperature of more than 900"C without having allowed the temperature of the slab to fall below the Ar3 point from the slab-forming step to the rolling step. In the present invention, a subsequent annealing can be carried out by a box annealing process or by a continuous annealing process. When a continuous annealing process is adopted, a higher coiling temperature such as from 650 to 750"C after hot rolling is favourable to develop the drawability of sheets.

It will thus be appreciated that when using a continuously cast steel slab or a bloomed steel slab made of Al-killed steel to perform this invention, AIN precipitation does not take place between the high temperature formation of the slab and the hot rolling, because the slab temperature is kept above the relatively low temperature of about 900"C. The method of this invention can thus obtain the best effects of these precipitates in the successive steps of forming the steel into the final product.

When a soft material such as the abovedescribed Al-killed steel is used according to the method of the present invention to produce a hot rolled steel for cold rolling into cold rolled steel sheet, the steel composition should be limited as follows (by weight): C 10.15% Mn < 0.50% N = 0.0020 - 0.0150% Sol Al = 0.015 - 0.10% with the balance being Fe and impurities.

Carbon must not be present in an amount of more than 0.15% because when more than 0.15% carbon is present, hardening of the hot and subsequently cold rolled steel sheet occurs, and the workability of the steel is also reduced.

Manganese must also not be present in an amount of more than 0.50% in order to ensure good workability, because when the manganese content increases to more than 0.50%, extreme deterioration of the workability is caused.

Furthermore, in Al-killed steel, it is necessary to develop a recrystallized texture in which the [111] planes of crystals are parallel to a rolling plane, so as to increase the workability of the steel making it suitable for use as a deep drawing cold rolled steel sheet. For this reason, the amount of soluble aluminium and the amount of nitrogen must be kept within the ranges of 0.015 to 0.10% Sol Al and 0.0020 to 0.015% N, respectively. If the amounts of these elements are kept within these ranges, the Al-killed steel can be formed into a hot rolled steel strip or sheet having excellent workability.

The invention can also be used for making high-strength hot rolled Al-Si-killed steel sheet, and here it is important to control the grain structure by the use of Al and N for producing a fine grain steel having excellent toughness. It is well known that in the usual prior art process for producing this kind of fine grain steel, the Al and N must be dissolved in the steel by reheating the slab immediately before the hot rolling, and the hot rolled strip must be subjected to a final hot rolling at a temperature above the Ar3 point. This strip is then coiled at a relatively low temperature, e.g. from 500 to 650"C, to keep the Al and N dissolved or in the precipitation state of AIN, so that a fine grain will be produced in a subsequent step, such as a normalizing step for precipitating the AIN. After a detailed study of AIN precipitates in Al-Si-killed steel, it has been found that in order to obtain a high-strength Al Si-killed steel having excellent workability and toughness, the cast slab or the bloomed slab must be directly fed to the hot rolling mill without allowing the temperature to fall below the Ar3 point. If necessary, heat can be added to maintain the slab above the minimum temperature.

When Al-Si-killed steel is to be formed into a high strength steel according to this invention, the starting composition should be limited as follows (by weight): C s 0.21% Mn = 0.70 - 1.60% Si = 0.10 - 0.40% Sol Al = 0.015 - 0.10% N = 0.0015 -- 0.0150% with the balance being Fe and impurities.

Carbon is effective for increasing the strength, but an excessive amount of C causes a deterioration in the toughness and welability of the steel, so that the carbon content must be limited to not more than 0.21%. Manganese and silicon are also effective to ensure that the strength is good without causing a deterioration in the toughness, but excessive amounts of these elements cause a deterioration in the weldability. For this reason, the manganese content must be limited to the range of 0.70 to 1.60% and the silicon content must be limited to the range of 0.10 to 0.40%. Al and N, which are used to obtain the fine crystal grain which gives the steel its good toughness, must be limited to the range of 0.015 to 0.10% for Sol Al and the range of 0.0015 to 0.0150% for N. If the Sol Al and N are kept within these ranges, the Si-AI killed steel which is subjected to the method of the invention to produce hot rolled and, if desired, normalized steel will have good toughness qualities. In addition, a Si-AI killed steel produced according to the present method is a weldable steel material having excellent toughness.

In producing a high strength steel which includes Ti, V and/or Nb it is very important to increase the strength by providing in such steel finely and uniformly dispersed precipitates of TiC, V(CN) and/or Nb(CN).

For this purpose, Ti, V and/or Nb and C and N must completely be dissolved in the hot slab before the hot rolling operation and then after the finishing of the hot rolling, TiC, V(CN) and/or Nb(CN) must be precipitated in the hot rolled strip.

After a careful study of the precipitation of TiC, V(CN) and/or Nb(CN) in steel which contains Ti, V and/or Nb it has been found that to obtain a desired high strength in the hot rolled steel, a cast slab or a bloomed high temperature slab must be directly fed to the hot rolling mill without allowing the slab temperature to fall below the Ar3 point.

If necessary, heat can be added to maintain the slab above the minimum temperature.

For a high strength steel which contains Ti, V and/or Nb and is to have a tensile strength between 50 kg/mm2 and 70 kg/mm2, the starting steel composition should be limited as follows (by weight): C = 0.06 - 0.20% Mn = 0.50 - 2.0% Si = 0.30 - 0.5% At least one of Ti, V or Nb, with the limitations: Ti = 0.01 N 0.10% V = 0.01 N 0.15% Nb = 0.01 - 0.10% with the balance being Fe and impurities.

Carbon, manganese and silicon are basic elements for ensuring good workability and achieving the desired strength level, and for these reasons, these basic elements must be present in minimum amounts of more than 0.06% carbon, more than 0.50% manganese and more than 0.03% silicon, respectively.

However, excessive amounts of these elements cause a loss in the desired workability which is required for hot-rolled strength steel. Therefore, the maximum amounts of these elements must be limited to not more than 0.20% carbon, not more than 2.00% manganese, and not more than 0.50% silicon, respectively.

Concerning the additional elements Ti, V and/or Nb, these must be present in amounts of 0.01 to 0.10% Ti, 0.01 to 0.15% V, and 0.01 to 0.10% Nb respectively. If these additional elements are present in smaller amounts than the above described amounts, they will not have sufficient influence to increase the strength. On the other hand, if amounts in excess of those set forth are added, no further effect is achieved. These additional elements can be added to the steel singly, or two or more may be added, according to the required strength and toughness.

The steel can be modified by including other elements such as P, Ni, Cr, Mo, Cu and Al, which increase the degree of corrosion resistance, wear resistance or other properties. If the steel is to have increased strength, the maximum amount of these elements which can be added without reducing the effect of the added Ti and/or V is about 1%.

This invention also extends to steel sheet whenever made by a method as described above of this invention.

The invention will now be described in greater detail, referring to the accompanying drawing which is a graph of the relationship between the tensile strength of the finished steel and the minimum slab temperature before hot rolling.

The drawing shows the manner in which the lowest temperature of a cast slab before heating or hot rolling influences the strength of the final steel, both with and without Nb.

From the drawing it can be seen that the critical feature affecting the strength of Nb containing steel formed from a slab rolled at a temperature of 1050"C is the minimum temperature to which the slab has been allowed to fall prior to the hot rolling. When the minimum temperature of the cast slab is above Ar3 point before reheating to 10500C for rolling the strength of the steel is kept high. The lowest temperature to which a Nb containing steel slab can be allowed to fall is about 800"C. The heating necessary to raise the temperature from about 800"C to the slab rolling temperature is not a reheating in the conventional sense, but rather is only heating to maintain a high temperature or to adjust the temperature.

The strength phenomenon is due to the way in which the precipitate, Nb(CN) is formed in steel which has not been allowed to cool below the Ar3 temperature between slab forming and hot rolling. When the method of the present invention is used and the slab temperature is not allowed to fall below the Ar3 point, the precipitation of Nb(CN) does not occur before final finishing hot rolling, and this Nb(CN) is finely precipitated only after the final finishing hot rolling. This results in an increase in the strength of the steel.

On the other hand, if the temperature of the slab is allowed to fall below the Ar3 point, Nb(CN) is completely precipitated, and it is not completely redissolved and uniformly dispersed in the steel even if the slab is reheated up to 10500C. Therefore, the precipitation of Nb(CN) before hot rolling is detrimental to the final strength.

As can be seen from the drawing, if the temperature of the slab is not allowed to fall below the Ar3 point so that a reheating step is unnecessary, the strength of the final product produced from the slab by hot rollling is as high as in the case of a minimum slab temperature of about 1000 C.

After hot rolling, the rolled steel containing Ti, V and/or Nb, is coiled at a low temperature, e.g. from 450 to 6500C, to cause precipitation of the TiC, V(CN) and/or Nb(CN) and is then subjected to cold rolling and box or continuous annealing for obtaining a high strength cold rolled steel sheet having excellent workability. The high strength is produced by uniform dispersion of the carbide and nitride precipitates as above described.

The composition of the steel employed in producing the accompanying drawing is as follows: from 0.10 to (3.12%C; from 0.21 to 0.24% Si; from 1.25 to 1.35% Mn; and (for some cases) from 0.04 to 0.05% Nb; with the balance being usual steel-making ingredients, unavoidable impurities and iron.

Certain specific examples of the invention will now be set out.

Example I.

Al-killed steels having slightly different compositions as shown in Table 1 were produced in a converter and possibly treated by a vacuum degassing treatment, the steels then being formed into a slab either by a continuous casting process or by a blooming operation after being cast in an ingot-making process. The thus-formed slabs were directly hot rolled, though if necessary temperature maintaining heating or reheating was used, and hot rolled according to the conditions in Table 1 to obtain a hot rolled steel having a thickness of 2.8 mm. The thus-obtained Al-killed hot rolled steel sheet was further subjected to a cold rolling step to obtain a final thickness of 1.0 mm, after pickling. Thereafter, a recrystallization annealing at 7100C for 6 hours was carried out and the steel sheet was finally temper-rolled to reduce the thickness by about 1.2%.

Table 1 shows the specific chemical compositions of the steels processed according to the present invention and the mechanical properties of steel sheets resulting from the processing steps. For the steels A-l to Ad, the respective slabs were not allowed to fall below a temperature of 900"C, i.e. not below the Ar3 point. In some cases, the temperature was maintained or increased slightly to the temperature at the time of charging into the rolling mill for hot rolling. The steel A-7 was directly hot rolled into a hot strip without maintenance heating and without the temperature of the slab falling below the Ar3 point from the time of blooming or continuous casting to the time of hot rolling.

On the other hand, the slabs formed from steels B-I to B-3 were allowed to fall below 850"C, i.e. below the Ar3 point, before being charged into a reheating furnace for heating to 1100"C for hot rolling.

A comparison of the qualities of the steels A I to A-7, produced according to the present invention, with the qualities of steels B-I to B-3 shows that the final product of the steels processed by the method of this invention is much softer, has a lower yield point and lower tensile strength, and also has a greater elongation. In addition, the steels A-I to A-7 have excellent properties, such as a high Er value and a high revalue and also have excellent deep drawability and stretchability. In the steels B-I to B-3, the slabs were allowed to fall to a temperature below the Ar3 point and AIN was precipitated at the time of this cooling, so that AIN was not completely dissolved and uniformly distributed in the steel when the slabs were reheated prior to hot rolling. Therefore, the T value of these products was very low. Steels A-I to A--5 are particularly to be noted, since the slab temperature at no time was lower below the Ar3 point and no AIN precipitation occurred in the slab prior to the end of rolling, even when the slab was heated up to 1100"C at the start of rolling. As a result, it was possible to obtain steel sheets having a high r value, i.e. more than 1.6, and a high Er value, i.e. more than 12.0. It is to be noted that the steel B--4, produced by a conventional process which involved reheating the cold slab to 12500C to dissolve the precipitated AIN, then subjecting the reheated slab to the usual hot rolling and cold rolling steps, is poor in that the steel has a low yield point, as well as a low Er value and a low r value.

From a theoretical standpoint, concerning the steels A-I to A-7, there are two major factors.

Firstly, AIN is not precipitated before the hot rolling operation.

Secondly, Al and N are uniformly dispersed and dissolved in the entire high temperature slab after blooming or casting and solidification and the precipitation of AIN starts for the first time at the time of recrystallization annealing. As a consequence, a good recrystallization texture develops which gives the steel good workability.

In the steel B--4, AIN is completely precipitated in the slab during the slab cooling step before hot rolling, and although the AIN is redissolved to Al and N during the reheating step, it is not uniformly dispersed in the slab because of the limited conditions during the actual heating operation, such as the heating time and temperature, and it is difficult to develop a preferably recrystallization texture for obtaining good workability by the subsequent recrystallization annealing.

Example II.

Molten Si-AI killed steel having a ladle composition of 0.15% C, 0.25% Si, 1.35% Mn, 0.013% P, 0.014% S, 0.03% Sol Al, 0.0045% N with the balance being Fe and impurities, was prepared in a 100-ton converter and cast into a slab by a continuous casting process. Al-Si-killed steel slabs obtained in this way were treated according to the conditions shown in Table 2. Each slab was hot rolled to a thickness of 25 mm and air-cooled, and the mechanical properties were determined. Furthermore, the hot-rolled steel was then annealed at 890 C for 15 minutes and the mechanical properties determined.

According to the test results, the steel of coils C-I and C-2 had better properties, such as yield point, tensile strength, elongation and charpy value, than the steel of coils D-l and D-2, whch were produced by a conventional process. In the steel or coils C-I and C-2, which were directly hot rolled without allowing the temperature to fall below the Ar3 point before the hot rolling step, Al and N were caused to precipitate after the hot rolling stage in fine grains to form a finely grained steel structure in which the precipitated aluminum nitrides are distributed uniformly throughout the steel structure. Such fine grain steel is characterized by having excellent strength and Charpy values, as seen in Table 2.

On the other hand, in the steel of coils D-I and D-2, the temperature of which was allowed to fall below the Ar3 point prior to hot rolling, AIN was completely precipitated when the slab cooled, and complete redissolution of the AIN was not achieved by the relatively low reheating temperature of coil D-l. In coil D-2, in which the AIN was dissolved at the high reheating temperature, Al and N were not dissolved and dispersed uniformly throughout the entire slab. Therefore, the benefits of the Al and N were not fully achieved in the steels of coils D-l and D-2. After the hot rolled and air-cooled coils C-I, C-2, D-l and D-2 were annealed by heating at 8900C for 15 minutes, and then air cooled, the mechanical properties of the steel of coils C-I and C-2, such as yield point, tensile strength, elongation. Charpy value and grain size, were good as compared with the same properties of the steel of coils D-l and D-2.

Example Ill.

Ti, V and Nb containing steels having the composition shown in Table 3 were cast into slabs having a temperature of more than 750 C. Slabs having compositions E-1 to E-6 were directly hot rolled, though some heating to adjust the temperature may have been applied. The slabs having compositions F-1 and F-2 were air cooled to the ambient temperature and then reheated and hot rolled. From the results of tests to determine the mechanical properties, as shown in Table 3, the finished steel from compositions E-1 to E-6 had higher values of tensile strength and toughness (VE60) as compared to the steel of compositions F-I and F-2. Even though the steel compositions E6 and F-I are the same, the treatment of steel F-l of cooling to a temperature below the Ar3 point prior to rolling gave a lower strength than steel E6, the temperature of which latter was not allowed to fall below the Ar3 point prior to hot rolling.

Example IV.

Ti, V and Nb containing steels having the compositions G-l and G-2 were cast and some of the cast slabs whilst at a temperature of more than 800 C were directly charged to a heating furnace, and then hot rolled without allowing the temperature of the slabs to fall. Other slabs from similar steels having compositions H-I and H-2 were similarly cast but then cooled down to ambient temperature, then reheated and hot rolled. The thus-obtained hot rolled steel strips having a thickness of 3.0 mm, were cold rolled to a thickness of 1.0 mm, subjected to annealing at 7000C for 2 hours and then temper rolled at a reduction rate of 1.5%. Th

TABLE 1 Treatment of Al-Killed Steel

I 1 00 c N O rC, In v, s \b oO ri 9 iC m - p9q v m = N n N n (s n bi ~ < ~ Hot rolling a X ~ t~ N ?} F - oo m t ~ ;i 8 F v D X uD s m t \0 Y 0 c 60 g: p%uals ~ E do es o ~- A xD > ~ do es. "I u a Q E t4 E Chemical Composition* rr m e m m lu!od V, cu u 03 q I? \s oo pla! C E d a, a; v; r- y, o c m r VE ci - N cl %u!lIO' uo!l3npai v MV1 0.Q55 0.27 0.01 Q ou 0.078 0.0100 oN 1000 om 490 m o oN o o Sullloo ~ or A-S duual 0.1:9 0.01 0.01 0.01 0.035 0.Q050 985 1050 m 64 15.2 o 48.2 12.9 1.89 :C A-6 0.060 0.3:0 0.01 0.01 OQ 00 0.0045 950 1250 m 00 00 Sulllol 0.29 0.01 0.01 0.01 0.060 0.0055 bo 1050 900 o o o 64 0 000v, B-1 0.043 ri X S 1 o 1100 875 o N o 64 < H N drual qals amleladuual o 0.Q49 0.29 ur 0.01 0.Q1 0.048 0.0048 o 1100 875 o o 22.5 36.1 44.8 11.8 1.35 qBlS UmW!U!UI 0.29 ofi o > 0.Q57 0.0076 1100 o o 520 6 23.6- 34.8 43.5 11.3 1.35 N m 0.050 ur O o vo 0.01 0.01 0.050 0.0045 m 1250 oo 550 m 17.8 31.2 46.4 11.9 1.55 b e \0 0 v) m t F t Z o o o ~ o o o o o o o o, o ob o o o o o. o o. o, ~ o o o o o o o o o o o ~ oo vn 00 v) o o 3s oo > o Zo o o o o o t- o o, o. o- o g m o o o o o o o o o o o .~ ut o o o o o o o o o o o .W O O O O O O O O O o X o.- o o o- o o o. o o o o U o o o o o o o o o o o Cd H H ~ ~ ~ H ~ H ~ H .~ .~ o o o o o o o o- o o o oE m O O o O c O s cw O s O U = t o afi o o os F 2 a- > - es- A - > *n o o o o o o o o o o o H 0O v) o o cs n o U) t t m ~ a0 m t t Ví m U o- o o o o o o- o, o, o o o o o o o o o o o o o uz < < < : < < < < &verbar; &verbar; m :: UO!}UXAUI luasald lusugeall I Sulplovau leuollua,suoD lUauuleal,L *Balance of all compositions is Fe and impurities.

TABLE 2 Treatment of Al-Si-Killed Steel

id V! 9 9 OZ r- r- Y, C 600 oin C c' W~ es o, o o Mechanical Properties after annealing > \u4NE Properties (as hot rolled) (8900C x o, minutes) slab hot ~ E m m m rci slab at start finishing point strength Y.P. T.S.

Coil temperature of rolling temp. (kg/ (kg/ El vE-20 CS (kg7 (kg/ El vE-20 CS sci ut O 1040 1040 a 36.6 53.4 29.0 15.0 6.5 m Qn , oo'oo' 'I: hln! IN O C: M P; r; r; r; Cd Uz O 700 1100 m 34.2 52.0 28.5 13.2 5.7 m = 20 1250 900 35.5 | 14.5 Ng s XD t U^l zc~ E e e < eo Wo5 É $ < > 8 8 sd o ca s ED < D O 0e oN o zz U U a a uotluaaul luasasd 01 Sulplovae luawleall luåwleal,l, leuolluaauoD TABLE 3 Treatment of Nb, Ti and V containing steel

"mOi V1. 00. g ;1 C? 9 CI C \0 00 f P1 WC 5:z < Cu 8 8 ~ ~ X ~y S properties S Chemical Composition* ux o o o o o o o o Coil u ness Ga m No. dwal qelS ~ Mo " " " " " " E-1 0.1-2 0.24 1.25 0.06 0.01 0.02 0.03 20 1000 1250 65 19.5 S E-2 0.06 0.21 o 0.6:4 0.1.0 e 0.20 0.05 0.02 0.03 o oo oN 1050 61 12.8 C urnuIruruI . 0.1-0 3 1.31 0.07 0.01 0.04 o s ç s ç o " ao.S > 1.26 > 0.03 0.01 " ~ ~ ~ > . 0.07 o.- Oo. O.o, 0.09 0.20 0.01 0.Q3 0.03 N -cN ZO 6 60 00 a" o" 0.50 0.09 0.20 0.06 O,,01 0.03 0,02 1260 ;0 11010 5566 2J07 O 0.02 ,o O O 5; N hl N o'o'o'o' of composition -r: a ~ o o o' o o o' ~ o Q o'o' o' o > , om o E U * O O m t > ~ O O O O O 0 O O U n O O A O O O O O O O O U X l l uolluaau luasasd oi luawleasl Su!plosse luswlealxL leuo!luaauoD TABLE 4 Treatment of Nb, Ti and V containing steel

t 4 Ch.

Ei \d n \o r c3 N o O ~ É xo OO mi e 'e 't: mF h! q i? f" E C O ,E vE b b b d!lls pallo' o o Pl03 Jo É ~ ~ ~ S 0 ssauaa!Ul v Properties drlls Composition* J É '=0 ='=0 Y.P. T.S. El oy Coil (kg,/ (kg' 9' E No. C Si Mn Ni Mo Cu V Ti Al Nb (0C) (0C) (mm) (mm) mm mm su!llol Jo 0.31 0.07 0.Ql 0.04 850 1250 o 1.0 o I1?S It? U i= O o" duIa qtls Bu!eay aioJaq ^ o o dural qels cu N UIIIUIUIIIl V H-i 0.07 0.26 1.25 0.50 0.Q9 0.20 0.Q6 0.Q1 0.03 0.02 20 1050 3.Q 1.0 43.7 52.3 27. & z ~ 0.09 0.25 1.33 0.02 0.03'. 0.04 9 1250 0^ oo o o H o ' O'' of o' a ~ 00' O' U z wn E ~ o ~ H ~ Lo wi ç rx c ' H I l l vo c) Z; n7 c) X = o UOI}U2AU! luasasd Y o1 su!plovae 1uawleall D luswleal,L leuo!luaauoD *

Claims (12)

WHAT WE CLAIM IS:

1. A method for producing a rolled, lowcarbon steel material, comprising forming a steel slab from a low-carbon steel composition containing (by weight) at least one carbide or nitride forming element selected from 0.015 to 0.10% acid-soluble aluminium, 0.01 to 0.10% titanium, 0.01 to 0.15% vanadium, and 0.01 to 0.10% niobium, by a forming process in which the finished slab is at a temperature above the Ar3 point of the steel and in which the or each said element is dissolved and dispersed evenly throughout the slab, maintaining the temperature of the thus-formed slab at a temperature above the Ar3 point of the lowcarbon steel from the time of the formation of the slab until the commencement of rolling of the slab, and directly carrying out hot rolling of the slab while the temperature of the slab is above the Ar3 point of the low carbon steel.

2. A method according to Claim 1, in which the hot rolled steel material is cold rolled and then annealed to obtain finished steel sheet or strip.

3. A method according to Claim I or Claim 2, in which the starting steel composition is a Al-killed steel consisting of C in an amount not more than 0.15%, Mn in an amount not more than 0.50%, 0.0020 to 0.015% N, and 0.015 to 0.10% Sol Al, with the balance being essentially iron and unavoidable impurities.

4. A method according to Claim 3, in which the temperature of the slab is maintained at more than 900"C prior to the hot rolling.

5. A method according to Claim 1 or Claim 2, in which the starting steel composition is a Si-Al-killed steel consisting of C in an amount not more than 0.21%, 0.70 to 1.60% Mn, 0.10 to 0.40% Si, 0.0015 to 0.015% N and 0.015 to 0.10% Sol Al, with the balance being essentially iron and unavoidable impurities.

6. A method according to Claim 5, in which the hot rolled steel material is hotcoiled at a temperature in the range of from 500 to 6500 C.

7. A method according to Claim I or Claim 2, in which the starting steel composition consists of 0.01 to 0.20% C, 0.50 to 2.00% Mn, 0.03 to 0.50% Si, 0.0015 to 0.015 ó N and at least one of Sol Al in an amount of 0.015 to 0.10%, Ti in an amount of 0.01 to 0.10%, and V in an amount of 0.01 to 0.15%, with the balance being essentially iron and unavoidable impurities.

8. A method according to Claim 7, in which the steel starting composition is modified by additionally comprising at least one alloying element selected from P, Ni, Cr, Mo, Cu and Al.

9. A method according to Claim 8, in which the or each said additional alloying element is present in an amount not exceeding 1%.

10. A method according to any of the preceding Claims and substantially as hereinbefore described.

Il. A method according to Claim I, and substantially as hereinbefore described in Example I (steels A-I to As?), or in Example II (steels C-I and C-2), or in Example III (steels E-l to E--6), or in Example IV (steels G-l and G-2).

12. Hot rolled steel material whenever produced in accordance with a method as claimed in any of Claims I to II.