EP2729590B1

EP2729590B1 - Hot-rolled high-strength steel strip with improved haz-softening resistance and method of producing said steel

Info

Publication number: EP2729590B1
Application number: EP12735134.4A
Authority: EP
Inventors: David Crowther; Willem Maarten VAN HAAFTEN
Original assignee: Tata Steel Ijmuiden BV
Current assignee: Tata Steel Ijmuiden BV
Priority date: 2011-07-10
Filing date: 2012-07-10
Publication date: 2015-10-28
Anticipated expiration: 2032-07-10
Also published as: CN103649355A; EP2729590A1; CN103649355B; WO2013007729A1

Description

The invention relates to a hot-rolled high-strength steel strip with improved HAZ-softening resistance and method of producing said steel
Traditionally steels with a high yield strength have been made by austenitising and quenching, but by this technique, for instance optimal surface quality and impact toughness may not be achieved. Manufacturing expenses have also been high.
EP1375694 describes a method for producing a high-strength, high-toughness steel with good workability and weldability by means of hot rolling. However, the inventors have found that these steels are susceptible to HAZ-softening. The heat-affected zone (HAZ) is the area of base material which has had its microstructure and properties altered by welding or heat intensive cutting operations. The heat from the welding process and subsequent recooling causes this change in the area surrounding the weld. The heat-affected zone which forms adjacent to the weld in steels is one of the most common regions of weld failure. Modern thermo-mechanically rolled (TMCP) steels, because of their lean chemistry, and in particular the low carbon content, require carefully controlled welding parameters in order to achieve adequate strength in the weld HAZ. For conventional low-carbon steels and some TMCP pipeline steels such as X70 and X80, HAZ softening is often the result of high heat-input welding procedure because of the slow heat dissipation in the HAZ. For higher grade TMCP steels such as X100, HAZ softening can happen even under moderate welding heat-input because of the base metal's ultra fine grain size and its bainite- and martensite-dominated microstructure. This reduction in hardness and strength in HAZ makes it a weak point in a welded pipeline structure. Consequently it is important to limit the degree to which this softening takes place, given the nature of the base metal and the welding conditions. The areas which have received most study in regards to HAZ softening are high strength plate steels e.g. for linepipe applications and, more recently, AHSS for automotive applications.
WO2007/129676 relates to a hot pressed steel member made from a high carbon cold-rolled steel sheet which is austenitised, hot-pressed to produce a steel member and quenched to achieve a minimum tensile strength of 1.8 GPa. The hot-rolled steel that is to be cold rolled contains at least 50% ferrite.
EP2028284 discloses a seamless steel pipe and EP1662014 discloses a low carbon hot rolled steel plate which is coiled between 450 to 650°C and subsequently promptly reheated to between 550 and 750°C to achieve a three-phase ferrite, bainite and island martensite structure.
The object of the present invention is to achieve a high-strength hot-rolled steel strip that is less susceptible to HAZ-softening than the currently available high strength hot-rolled steel strip.
In a first aspect this object is reached by a hot-rolled high-strength micro-alloyed steel strip having a thickness of between 2 and 16 mm with improved HAZ-softening resistance having a microstructure comprising martensite, tempered martensite and/or bainite, and where the steel contains, in percentages by weight:

0.07 - 0.27% C;
0.8 - 2.0% Mn;
0.01-0.08% Al_sol;
0.2 - 1.5% Cr;
0.1-0.7% Mo;
0.0005-0.005 B;
0.01-0.07% Nb;
at most 0.5% Si;
at most 0.03% P;
at most 0.015% S;
at most 0.05% Ti;
at most 0.1% V;
at most 0.2% Cu;
at most 0.2% Ni;
at most 0.008% N;
optionally calcium additions for sulphide shape control, at most 0.015%;
other elements in amounts of impurity level, balance iron;

All compositional percentages are given as weight percent, unless indicated otherwise.
The inventive idea is based on the fact that by selecting the combination of chemical elements in the amounts prescribed and in particular the presence of niobium and molybdenum good hardening can be maintained. The steel structure is not critical for the segregation of manganese and carbon during the casting process owing to the low manganese and carbon content. The steel properties are not critical for local fluctuations of the coiling temperature in the strip, which facilitates the steel production and has an advantageous effect in the homogeneity of its mechanical properties, which again has a positive influence both in the flatness of the end product and in the residual stress. The steel sheet is highly suitable for welding and laser cutting, and at the same time it has a good fatigue strength irrespective of said thermal treatments. Further, the steel sheet has excellent bending properties, a good impact toughness as well as a good resistance to softening in tempering. The presence of the precipitating elements in solution which are subsequently available for precipitation during the cooling of the HAZ after welding ensures a significant improvement of the HAZ-softening resistance.
The steel according to the invention can be thermally cut, for instance by laser, into precisely defined shapes. It has been observed that a remarkably smooth cutting surface is achieved in a laser cut object. On the other hand, it has been found out that the strength difference between the basic material and the soft zone created in the technical cutting process, which zone is located in the immediate vicinity of the hardened zone, is small or absent also as a result of the presence of the precipitating elements in solution which are subsequently available for precipitation during the cooling of after thermal cutting. These together have an advantageous affect in the fatigue strength. In addition, the relatively low carbon content reduces the peak hardness of the hardened zone, so that the cutting surface is not sensitive to embrittlement and cracking, neither in the working of the object nor in practical use.
Carbon is an important element in controlling the strength, but needs to be limited to some extent to give a good balance in toughness, weldability and formability. The C content is kept relatively low between 0.07 and 0.27% to achieve a good toughness (all percentages in compositions are expressed in weight percent). In combination with a low temperature coiling temperature the microstructures will contain martensite and/or bainite. The exact amounts strongly depend on the composition, cooling rate on the run-out table and the coiling temperature. At relatively low C contents, the Ms temperature will be quite high, so the martensite will auto temper to some extent. Depending on the strength requirement suitable carbon-windows were found. For a steel having a yield strength of at least 960 MPa the carbon content is preferably at least 0.07 and/or at most 0.13%. For a steel having a yield strength of at least 1100 MPa the carbon content is preferably at least 0.13 and/or at most 0.18%, and for a steel having a yield strength of at least 1300 MPa the carbon content is preferably at least 0.19%, more preferably at least between 0.23 and/or at most 0.27%.
A suitable maximum carbon content is 0.27%.
Elements like manganese, chromium, molybdenum and boron provide hardenability to promote the formation of bainite and/or martensite. The manganese content is limited to between 0.8 and 2.0%. When the manganese exceeds the upper boundary, the risk of segregation becomes significant and this may adversely affect the homogeneity of the microstructure. At levels below 0.8% the effect on hardenability is insufficient. A suitable minimum manganese level is 1.1%.
Boron is added to promote the hardenability. It is important to avoid the formation of boron nitrides as this will render the boron ineffective for the promotion of the hardenability. The role of titanium in the composition according to the invention is to protect the boron because Ti forms titanium nitrides and as a consequence no BN is formed. The amount of alloyed boron is at least 0.0005% B (i.e. 5 ppm) but no more than 0.005% B (i.e. 50 ppm) in order to reduce grain size and to increase the hardenability. If titanium is added as an alloying element, the amount of titanium is typically at least 0.01% Ti but no more than 0.05% in order to bind the nitrogen N and to prevent the creation of boron nitrides BN. An alternative within the scope of the invention is to use aluminium to bind the nitrogen and thereby protect the boron. In cases where the formation of TiN-particles is deemed undesirable, e.g. because of their effect on Charpy-toughness, protecting the free boron by aluminium may be the preferred option.
Niobium added in an amount of 0.01 - 0.07 partly precipitates in the austenite during hot-rolling, thereby contributing to the grain refinement of the final transformed microstructure by the retardation of the recrystallisation of austenite. In addition, the niobium remaining in solution at transformation has a powerful effect on reducing transformation temperatures, especially at faster cooling rates, so it is also beneficial for hardenability. At low coiling temperatures (<500°C) the contribution of niobium by precipitation strengthening is expected to be small because the temperature is too low for precipitation of fine NbC. A suitable Nb content is at least 0.02%, preferably at least 0.025%.
V added in an amount of less 0.1% has a similar but less powerful effect as Nb in this case. However, the main reason for the addition of Nb and V is to improve the HAZ-softening resistance. In relevant parts of the HAZ, the thermal cycle is such that temperatures are reached which will allow precipitation strengthening by Nb and V, thus causing an increase in hardness as a result of the precipitation of elements which were kept in solution by the low coiling temperature. The major contribution is believed to be made by Nb and V carbides, nitrides or carbo-nitrides. To a lesser extent it is believed that MoC precipitates may form having a similar effect. If present, a suitable minimum V-content is 0.04%.
The steels according to the invention are aluminium-killed or aluminium-silicon killed steels in order to reduce the oxygen content to a minimum so that no reaction occurs between carbon and oxygen during solidification. The amounts of aluminium added to the steel during production therefore include those needed for deoxidation. The remaining amount in the end product, also called soluble aluminium (Al_sol) is between 0.01 and 0.08% Al. In the context of this invention the aluminium content referred to is soluble aluminium.
Silicon may also serve as a deoxidant in the steel of the present invention in addition to aluminium. It also acts as a solid solution hardener starting from about at least 0.10% Si and up to 0.50 % Si, which has an advantageous effect on the impact toughness and workability. Above 0.5% the silicon adversely affects the surface quality of the steel to an unacceptable extent and the removal of the hot rolling scale by pickling becomes increasingly difficult with increasing silicon content.
Phosphorus P contained as an impurity should be at most 0.03%, and sulphur S should be even lower and should be limited to at most 0.015%, which means that these contents are restricted in order to achieve good impact toughness and bendability. When necessary, further properties can be improved by treating the melt with cored wire containing Ca-Si or Ca-Fe(Ni). During calcium treatment, the alumina and silica inclusions are converted to molten calcium aluminates and silicate which are globular in shape because of the surface tension effect. The calcium aluminates inclusions retained in liquid steel suppress the formation of MnS stringers during solidification of steel. This change in the composition and mode of precipitation of sulphide inclusion during solidification of steel is known as sulphide morphology or sulphide shape control. This results in less nozzle clogging during casting, better mechanical properties as the long stretched MnS stringers act as crack initiating points. Typical amounts of calcium in the steel for sulphide shape control is 0.0015 to 0.015% Ca. A suitable maximum is 0.005% Ca.
Chromium should be between 0.2 and 1.5%. Molybdenum should preferably be between 0.1% Mo and 0.7% Mo. Both elements are added in order to increase hardening and tempering resistance. This enables precipitation at higher coiling temperatures, which can be used for decreasing and even preventing the softening of the steel, as well as for alleviating strength fluctuations caused by local temperature differences during the cooling of the coil. A suitable minimum molybdenum content is 0.15%.
Alloying with elements like copper and nickel, often used in steels of this strength level are preferably avoided in view of the surface issues associated with copper. As copper is often alloyed in conjunction with nickel to alleviate the adverse effects of copper this is also not needed. So nickel and/or copper are preferably present at most at impurity level or more preferably completely absent.
The microstructure of the steels according to the invention are characterised as a microstructure that consists predominantly of tempered martensite, characterised by small carbides in a Widmanstätten pattern, and/or bainite. Ideally the microstructure is free from ferrite and pearlite constituents as these will deleteriously affect the strength level to be reached. In practice it may be unavoidable that some minor patches of ferrite are present, but the amount may not exceed the level where the strength levels is significantly affected. The abovementioned deleterious and subsequently undesirable ferrite constituents which form at high transformation temperatures should be clearly distinguished from the ferritic part of bainite or Widmanstätten ferrite or acicular ferrite which form at low transformation temperatures. The former constituents are undesirable, the latter are not.
The hot-rolled steel strip according to the invention that is directly hot-rolled to the thickness 2 mm - 16 mm can be manufactured as wear-resistant and with different minimum yield strength. Typical threshold values in the marketplace are 960, 1100 and 1300 MPa, only by changing the analysis and/or the post-rolling cooling rate of the strip, and/or temperature before the coiling, within the scope of the invention. This kind of high yield strength steel can also be used in targets where the structures require properties typically demanded of structural steel, such as good workability, weldability and impact toughness, which means that the hot-rolled steel strip according to the invention is feasible also as weldable structural steel. In the steel analysis to be explained in the specification below, all content percentages are percentages by weight, and the rest of the steel that is otherwise not defined is naturally iron, Fe, and unavoidable impurities. The value of 960 MPa can be reached even to thickness values of the hot strip of up to 16 mm. For 1100 and 1300 MPa the values can be reached for values of up to 12 mm. The higher thickness results in a somewhat lower cooling rate, and therefore in a less enriched austenite prior to the phase transformation. This leads to a somewhat lower strength level, meaning that with these lean compositions the S1100 or S1300 level cannot be obtained for thicknesses of the hot strip over 12 mm. Preferably the minimum thickness is 3 mm and/or the maximum thickness is 10 mm. It should be noted that the values of the strength as defined in this invention are measured in the longitudinal direction (i.e. the tensile specimen is taken in the longitudinal direction of the strip (the direction of movement through the rolling mill)). Values in the transverse direction (i.e. the tensile specimen is taken in the width direction of the strip) may be different from the values in the longitudinal direction, and are usually higher than those in the longitudinal direction for strength and lower for elongation.
In an embodiment the carbon content of the steel is between 0.07 and 0.13% and the yield strength is at least 960 MPa.
In an embodiment the carbon content of the steel is between 0.13 and 0.18% and the yield strength is at least 1100 MPa.
In an embodiment wherein the carbon content is at least 0.19%, preferably between 0.23 and 0.27%, and the yield strength is at least 1300 MPa.
A suitable maximum tensile strength of the hot-rolled steel according to the invention is 1700 MPa.
According to a second aspect the invention is embodied in a method for manufacturing a hot-rolled high-strength micro-alloyed steel strip having a thickness of between 2 and 16 mm with improved HAZ-softening resistance and a yield strength of at least 960 MPa having a microstructure comprising martensite, tempered martensite and/or bainite, and where the steel contains, in percentages by weight:

₃

finish rolling to a final thickness of from 2 to 16 mm
cooling the hot rolled strip within at most 10 seconds from the last hot rolling pass to a coiling temperature of between 20 and 500°C at a cooling rate sufficient to transform the rolled microstructure into a microstructure comprising martensite and/or bainite.

The hot-rolling process is the conventional hot-rolling process, either starting from a slab having a thickness of between 150 to 350 mm, i.e. conventionally continuously cast slabs, or below 150, i.e. thin slab casting or even strip casting. Finish rolling is preferably while the steel is still austenitic to provide the steel with a fine grain and thereby good impact toughness. The coiling temperature is preferably low to achieve the desired mechanical properties. Preferably the coiling temperature is below 400°C. By manufacturing this type of steel by fast cooling or quenching directly after hot rolling, an excellent impact toughness is obtained because the phase transformation into martensite and/or bainite takes place from a fine-grained, worked austenite. It also improves the surface quality because the primary scale is removed in a descaler prior to the rolling. Moreover, there is no need for a very expensive lengthy additional annealing treatment to dissolve all precipitates. In a hot strip rolling line, the slabs to be rolled are reheated in a reheating furnace to a temperature range between 1100 to 1250°C, and held for several hours. In that case the dissolution of special carbides, such as Cr and Mo carbides, and the homogenization of the structure is as complete as possible. On the other hand, the growing of the austenite grain at the high heating temperature does not make the end product more brittle, because the austenite grains are refined first by recrystallisation during high temperature rolling in the initial stages of the rolling process and by transformation of the deformed austenite grain formed as a result of the retardation of austenite transformation during thermo-mechanical rolling in the last stages of the hot rolling process. The heavily deformed austenite grains transform into a very fine transformation product during cooling on the run-out table. This results in a high yields stress, combined with an excellent impact toughness. Manufacturing costs and production time can be further reduced if a thin slab casting and direct rolling facility is used where the elements like Cr and Mo carbides have not even yet precipitated before the rolling starts.
According to the invention, steel is manufactured at a final rolling temperature at which the steel is still austenitic, i.e. above Ar3 to a final hot rolled thickness. The cooling of the strip begins no later than 10 seconds after the last hot rolling pass, and it is cooled sufficiently rapidly to allow the austenite to transform into a bainitic and/or martensitic microstructure, the cooling rate preferably being at least 30°C/s, down to a coiling temperature in the range 20°C - 500°C, preferably down to a coiling temperature in the range 20°C - 450°C. The obtained result is typically a nearly completely bainitic and/or martensitic microstructure, so that the bainite and/or martensite content preferably is at least 90 % by volume, preferably at least 95%. The microstructure is also preferably free from ferrite formed at high temperatures and free from pearlite constituents, thus rendering the microstructure substantially fully bainitic/martensitic, where Widmannstätten ferrite or acicular ferrite is considered to be a bainitic structure for this purpose. In the coiling temperature range of below 100°C the martensite is not tempered, although some auto-tempering may occur at these low temperatures for the low carbon grades, whereas when the coiling temperature is at least 100°C, the martensite is tempered. At temperatures above 200°C the martensite is tempered and the carbon precipitated.
In a preferable embodiment the coiling temperature is at most 450°C, more preferably 425°C, for steels containing a carbon content of at most 0.12% or at most 275°C, more preferably 250 or 225°C, for steels containing a carbon content of at least 0.13%.
In an embodiment the cooling rate during the accelerated cooling after hot rolling and before coiling is between 5 and 100°C/s.
In an embodiment the finish rolling temperature is above Ar₃ and also below 920°C, and preferably also below 900°C.
According to a third aspect the steel according to the invention is used in the production of a part for an automobile, a lorry, ship-building, construction work, heavy haul equipment, earth-moving equipment or mobile cranes.

The invention is now further explained by means of the following, non limiting examples. Table 1 shows compositions of steels for each of the categories S700, S960, S1100 and S1300.

Table 1. Composition of the steels (all in weight% x1000, except Band N (ppm). All steels are Ca-treated, imp=impurity level)

Type	ID	C	Si	Mn	Al	P	S	Cr	Mo	B	N	Ti	V	Nb
S960
High Mn	4587	100	250	1770	30	14	5	260	250	20	48	33	imp	39
High Mn	4588	91	240	1780	34	15	5	500	5	20	60	35	imp	39
Low Mn	4589	110	220	1210	31	15	4	490	510	20	50	30	imp	38
Low Mn	4590	110	240	1180	32	15	4	1030	5	20	80	32	imp	39
Low Si	4591	90	100	1600	30	14	5	500	250	20	50	30	imp	40
Low Si	m	94	108	1590	38	11	2	502	263	25	52	27	47	40
Base	4699	99	230	1490	33	15	4	510	250	30	60	29	imp	40
+V	4671	94	240	1510	33	15	3	530	250	25	60	31	51	39
-Nb	4670	97	230	1500	35	15	4	540	250	25	60	30	imp	imp
+Cr-Nb	4673	110	240	910	32	15	4	1130	300	30	60	29	imp	imp
S1100
B	4814	155	140	1460	23	15	5	490	240	30	60	26	imp	29
+C	4815	170	140	1460	22	15	5	490	240	30	60	29	imp	29
+Mo	4816	165	140	1470	22	15	5	490	510	30	60	30	imp	30
-Cr	4817	155	140	1460	21	15	5	260	240	30	60	29	12	19
+Mo-Cr	4818	155	140	1460	22	15	5	250	500	30	60	28	imp	29
-Ti+Al	4819	145	140	1460	59	15	5	450	240	30	60	6	imp	29
+V	4820	155	150	1460	23	15	5	490	240	20	60	29	50	29
-Ti+Al+V	4819*	145	140	1460	59	15	5	450	240	30	60	6	50	29
S1300
Base		250	100	1490	27	14	3	480	510	20	60	3	imp	29
+B		230	100	1510	27	14	4	490	510	16	60	2	imp	30
+Al		250	110	1500	64	14	4	480	510	18	60	2	imp	28
+Ti		270	110	1510	27	14	4	490	510	20	60	25	imp	30

SEM microstructures show that all hot-rolled steel samples have a microstructure that consists predominantly of tempered martensite, characterised by small carbides in a Widmanstatten pattern. The tempering of the martensite will have resulted due to the high M_s temperature of these relatively low C steels, and also due to the slow cooling to simulate coil cooling. The pattern of the carbides, showing several variants of an orientation relationship between the carbide and the matrix, is characteristic of tempered martensite. The prior austenite grain boundaries were also visible, and showed some elongation along the rolling direction, as was observed using optical metallography. The strength levels depend strongly on the coiling temperature (see figure 1 for the S1100 materials and figure 2 for the S1300 material) which is to be understood in view of the desired microstructure that consists predominantly of tempered martensite.
Figure 3 shows the effect of HAZ-softening on the Y-axis for steels (a, c) according to the invention and for comparative steels (b, d). It is clearly visible that the inventive steels outperform the comparative steels in terms of a reduced softening of the HAZ. Steels b and d are S960 -Nb and S960 +Cr-Nb respectively (see table 1) and a and c are S960 Base and S960 +V respectively.

Claims

Hot-rolled high-strength micro-alloyed steel strip having a thickness of between 2 and 16 mm with improved HAZ-softening resistance having a microstructure comprising martensite, tempered martensite and/or bainite, and where the steel contains, in percentages by weight:
• 0.07 - 0.27% C;

• 0.8 - 2.0% Mn;

• 0.01-0.08% Al_sol;

• 0.2 - 1.5% Cr;

• 0.1-0.7% Mo;

• 0.0005-0.005 B;

• 0.01-0.07% Nb;

• at most 0.5% Si;

• at most 0.03% P;

• at most 0.015% S;

• at most 0.05% Ti;

• at most 0.1% V;

• at most 0.2% Cu;

• at most 0.2% Ni;

• at most 0.008% N;

• optionally calcium additions for sulphide shape control, at most 0.015%;

• other elements in amounts of impurity level, balance iron;
the yield strength of the steel strip being at least 960 MPa.
Steel according to any one of the preceding claims having a chromium content of at least 0.2%.
Steel according to any one of the preceding claims having a nickel content at impurity level.
Steel according to any one of the preceding claims having a copper content at impurity level.
Steel according to any one of the preceding claims wherein the carbon content is between 0.07 and 0.13% and having a yield strength of at least 960 MPa.
Steel according to any one of claims 1 to 4 wherein the carbon content is between 0.13 and 0.18% and having a yield strength of at least 1100 MPa.
Steel according to any one of claims 1 to 4 wherein the carbon content is at least 0.19%, preferably between 0.23 and 0.27%, and having a yield strength of at least 1300 MPa.
A method for manufacturing a hot-rolled high-strength micro-alloyed steel strip having a thickness of between 2 and 16 mm with improved HAZ-softening resistance and a yield strength of at least 960 MPa having a microstructure comprising martensite, tempered martensite and/or bainite, and where the steel contains, in percentages by weight:
• 0.07 - 0.27% C;

• 0.8 - 2.0% Mn;

• 0.01-0.08% Al_sol;

• 0.2 - 1.5% Cr;

• 0.1-0.7% Mo;

• 0.0005-0.005 B;

• 0.01-0.07% Nb;

• at most 0.5% Si;

• at most 0.03% P;

• at most 0.015% S;

• at most 0.05% Ti;

• at most 0.1% V;

• at most 0.2% Cu;

• at most 0.2% Ni;

• at most 0.008% N;

• optionally calcium additions for sulphide shape control, at most 0.015%;

• other elements in amounts of impurity level, balance iron;
the strip being finish hot-rolled above the Ar₃-temperature, wherein the method includes at least the following steps:
• finish rolling to a final thickness of from 2 to 16 mm

• cooling the hot rolled strip within at most 10 seconds from the last hot rolling pass to a coiling temperature of between 20 and 500°C at a cooling rate sufficient to transform the rolled microstructure into a microstructure comprising martensite and/or bainite.
Method according to claim 8 wherein the coiling temperature is:
• at most 450°C for steels containing a carbon content of at most 0.12% or

• at most 275°C for steels containing a carbon content of at least 0.13%.
Method according to any one of claims 8 to 9 wherein the cooling rate during the accelerated cooling after hot rolling and before coiling is between 5 and 100°C/s.
Method according to any one of claims 8 to 10 wherein the finish rolling temperature is above Ar₃ and also below 920°C, and preferably also below 900°C.
Use of the steel according to claim 1 to 7 in the production of a part for an automobile, a lorry, heavy haul equipment, earth-moving equipment or mobile cranes.