Patent Application
20100294401
The present invention relates to a high strength bainitic steel, to a process for producing seamless pipes for OCTG applications and to the use of this steel for OCTG applications.
Quenched and tempered martensitic steels are currently broadly used to produce high strength seamless pipes for OCTG applications.
One interesting alternative to get improved mechanical properties is the use of carbide-free bainitic steels in the as rolled or as rolled and tempering conditions. The chemical composition of these steels must be carefully designed to suppress the ferrite and pearlite reactions during the slow air cooling from the austenitic range after hot rolling.
The loss of toughness and ductility commonly observed in bainitic steels is usually related to the presence of coarse cementite particles between the bainitic ferrite sheaves. In order to avoid this problem, it was proposed to inhibit the cementite formation by the addition of more than 1 wt % of Silicon or Aluminum. These elements can not be dissolved in cementite, and hence suppress its precipitation. From the document WO96/22396 there is known a carbide-free high Si/Al bainitic steel, but it is used for different applications than for OCTG applications. In particular WO 96/22396 discloses a method of producing a bainitic steel product, whose microstructure is essentially carbide-free, comprising the steps of: hot rolling the steel product and either cooling the steel from its rolling temperature to ambient temperature continuously and naturally in air or by continuously accelerated cooling. The cooling rates used are between 225 and 2° C./s, therefore comprising very high cooling rates.
The material is produced as rolled or after accelerated cooling, and the product is always intended for different applications than for OCTG applications.
It is a fact that bainitic steels in the as rolled condition or after accelerated cooling can not be directly used for high strength OCTG applications. Due to the low yield to tensile strength ratio, the required tensile and impact properties cannot be achieved, in particular for some OCTG applications.
The need is therefore felt to provide a steel composition and a process for producing seamless steel tubes having high strength and toughness, suitable for OCTG applications.
The main object of the present invention is to provide an improved process for producing seamless free-carbide bainitic steel tubes, having high strength and toughness, suitable for OCTG applications.
Another object of this invention is to provide a steel composition for producing high strength seamless tubes for OCTG applications, with high Yield Strength (YS) and good toughness.
The present invention, therefore, proposes to achieve the purposes described above providing a process for the production of high strength bainitic steel seamless pipes comprising the following steps:
a) providing a steel having a composition comprising 0.2-0.4% by weight of C 0.05-1.5% by weight of Mn; 1.0-2.0% by weight of Si and 0-0.5% by weight of Al or, alternatively, 1.0-2.0% by weight of Al and 0-0.5% by weight of Si; 0.5-2.0% by weight of Cr; 0.2-0.5% by weight of Mo; 0.5-3.7% by weight of Ni; the remainder being iron and inevitable impurities;
b) hot rolling said steel at a predetermined temperature such as to obtain a seam-less steel pipe;
c) continuously cooling the steel from the rolling temperature naturally in air or by a controlled cooling with an average cooling rate comprised between 0.10 and 1.0° C. per second in order to obtain mainly bainitic structures.
The product directly obtained by said process is a seamless steel pipe for OCTG applications that, according to claim 10, has a mainly cementite-free bainitic microstructure and displays a yield strength of at least 140 ksi and a Charpy V-notch impact energy at room temperature of at least 50 J (full size samples).
According another aspect of the invention, there are provided:
The core of the invention is to use a mainly cementite-free bainitic structure in seamless tubes for high strength OCTG applications.
Advantageously a low temperature tempering treatment in the steel of the invention is also a non-conventional treatment because it is not used to improve toughness, since Charpy results are only marginally improved by this treatment, instead it is aimed at increasing yield strength through precipitation of small transition carbides and dislocation pinning by interstitials.
The advantages ensuing to the steel of the invention are the improvement in strength-toughness over tempered martensitic steels, and the simplified thermal treatment, because only a low temperature tempering treatment is needed, without previous quenching.
In comparison to the quenched and tempered martensitic steels, carbide-free bainitic steels in the condition as rolled and with low temperature tempering have, therefore, the following two major advantages:
a. quenching is not necessary and by avoiding the quenching treatment the microstructure results far more homogeneous, which allows thick walled tubes to be produced;
b. for the same steel composition, in comparison to conventional tempered martensitic structures, a better combination of strength and toughness can be achieved, in particular by tempering as rolled carbide-free bainitic structures.
The foregoing and other objects will become more readily apparent by referring to the following detailed description and the appended drawings in which:
a show the microstructure of B2 as rolled and tempered at 300° C. (transmission electron image);
The steel of the invention has a composition in weight percent comprising:
C: 0.2-0.4; Mn: 0.05-1.5; Si: 1.0-2.0 and Al: 0-0.5 or, alternatively, Al: 1.0-2.0 and Si: 0-0.5; Cr: 0.5-2.0; Mo: 0.2-0.5; Ni: 0.5-3.70; S: 0-0.005; P: 0-0.015; Ca: 0-0.003; O: 0-0.005; Cu: 0-0.15; N: 0-0.01; balanced iron save for incidental impurities.
A first preferred composition of the steel comprises in weight percent:
C: 0.23-0.30; Mn: 0.05-1.0; Si: 1.2-1.65 and Al: 0-0.5 or, alternatively, Al: 1.2-1.65 and Si: 0-0.5; Cr: 0.7-1.8; Mo: 0.2-0.3; Ni: 0.5-3.6; S: 0-0.005; P: 0-0.015; Ca: 0-0.003; O: 0-0.002; Cu: 0-0.1; N: 0-0.01; balanced iron save for incidental impurities.
A further advantageous preferred composition of the steel comprises in weight percent:
C: 0.23-0.30; Mn: 0.05-0.7; Si: 1.2-1.6; Al: 0.01-0.04; Cr: 0.7-1.4; Mo: 0.2-0.3; Ni: 2.0-3.6; S: 0-0.003; P: 0-0.015; Ca: 0-0.002; O: 0-0.0015; N: 0-0.0080; Cu: 0-0.1; balanced iron save for incidental impurities.
The microstructure of the steel is essentially a fine cementite-free bainite with minor fractions of retained austenite and martensite. It is obtained after hot rolling and continuously cooling the steel from its rolling temperature naturally in air or by a controlled cooling.
Advantageously, the average cooling rate after hot rolling has to be in the range between 0.10 and 1.0° C./sec, preferably between 0.2 and 0.5° C./sec, in order to obtain mainly bainitic structures for the range of steel compositions tested. This is the case of tubes naturally cooled in air with wall thickness between 8 mm and 16-18 mm. For thicker or thinner tubes a controlled cooling with said average cooling rate may be needed to achieve the desired structure after hot rolling.
In spite of the high hardness, the as rolled bainitic structures have a low yield to tensile strength ratio, thereafter in this condition it is not possible to reach very high values of yield strength and at the same time the high impact properties needed for some OCTG applications, for example deep well applications. Advantageously, in order to meet these requirements a tempering treatment at low temperatures (200-350° C.) has to be performed. During this treatment the yield strength strongly increases due to transition carbide precipitation and dislocation pinning by interstitials; and the impact properties are not impaired. As a consequence, a good combination of strength and toughness (suitable for high strength OCTG applications) can be achieved. The duration of this tempering treatment is about 30-60 minutes.
Regarding steel chemistry, in order to minimize coarse cementite precipitation, detrimental to toughness, during continuous cooling from the hot rolling temperature and during tempering, high Si or Al contents are used.
Advantageously, 1-2 weight percent of Si or Al has to be used. Both elements have similar effects on carbide precipitation during the bainitic reaction, because of their low solubility in cementite. If high Si is used, the Al content of the steel will be lower than 0.5 weight percent. Conversely, if high Al is used, the Si content of the steel will be below 0.5 weight percent.
The intermediate carbon contents, preferably 0.23-0.30 wt %, have the function of depressing the bainitic start temperature and getting microstructural refinement. Moreover, in order to achieve high strength in the as rolled condition, the transformation temperature is deplected by Mn, Ni, Cr and/or Mo alloying additions.
In particular, in order to avoid ferrite and perlite formation during natural air cooling Ni+2Mn has to be between 2 and 3.9, where Ni and Mn are concentrations in weight percent. Fulfilling this condition, Ni can be partially replaced by Mn in the steel composition.
However, in a preferred embodiment of the composition, Ni-content is present at high concentrations, preferably 2.0-3.6 wt %, for improving toughness while Mn is kept as low as possible, preferably 0.05-0.7 wt %, in order to avoid the formation of large blocks of retained austenite.
Mo is added at the herein specified levels, preferably 0.2-0.3 wt %, to avoid P segregation to interphases at low temperature.
Cr is added at the herein specified levels, preferably 0.7-1.4 wt %, to avoid, together with Mo and Ni, the ferrite and perlite formation during air cooling and to improve microstructural refinement by lowering the bainitic start temperature.
O is an impurity present mostly in the form of oxides. As the oxygen content increases, impact properties are impaired. Accordingly, a lower oxygen content is preferred. The upper limit of the oxygen content is 0.0050 wt %; preferably below 0.0015 wt %.
Cu is not needed, but depending on the manufacturing process may be unavoidable. Thereafter, a maximum content of 0.15 wt % is specified.
The contents of unavoidable impurities such as S, P, Ca, N, and the like are preferably low. However, the features of the present invention are not impaired as long as their contents are as follows: S not greater than 0.005 wt %; P not greater than 0.015 wt %, Ca not greater than 0.003 wt % and N not greater than 0.01 wt %; preferably S not greater than 0.003 wt %; P not greater than 0.015 wt %, Ca not greater than 0.002 and N not greater than 0.008 wt %.
The following examples are useful for better defining the invention and to point out the influence of the chemical composition and of the process steps on the behavior of the steel. In particular, the feasibility of producing high strength bainitic steels that fulfill the tensile and impact requirements of deep well OCTG products is investigated.
The following tasks were performed:
1. Three alloys (B1, B2 and B3 steels) were designed. The materials were laboratory cast and hot rolled in a pilot mill.
2. The as rolled microstructures were studied under optical and scanning electron microscopes. X-ray diffractometry was used to quantify the amount of retained austenite.
3. Standard tensile and Charpy impact (at −20° C., 0° C. and room temperature) tests were carried out on as rolled samples. Hardness measurements were also performed.
4. The transformation behavior of the alloys was studied in a thermomechanical simulator. CCT diagrams were measured for all steels.
5. To determine the effect of different thermal cycles on mechanical properties, normalization and tempering treatments were performed on as rolled plates. Hardness, tensile and impact measurements were conducted on the heat treated samples.
The alloy design was aimed to produce a microstructure mainly composed of bainitic ferrite and films of retained austenite during air cooling from the austenitic range. From calculations performed with a computer program, it was estimated that, for tube thicknesses between 24 mm and 6 mm, the average cooling rate at the exit of the hot rolling mill (rolling temperature: 1100-950° C.) is in the range between 0.1° C./sec and 0.5° C./sec. Several chemistries were designed to get the desired microstructure during cooling at the above mentioned rates. The concentration of each element was selected with the aid of a metallurgical model for the prediction of TTT diagrams (H.K.D.H. Bhadeshia, “A thermodynamic analysis of isothermal transformation diagrams”, Metal Science, 16 (1982), pp. 159-165). The resulting chemistries (B1, B2 and B3) are shown in Table 1.
The only difference between B1 and B2 steels was the carbon content, which was changed in order to study its effect on microstructure and mechanical properties. In B3 steel several changes were performed in comparison with the previous alloys: C was increased to improve microstructural refinement and Si was replaced by Al as the element used to inhibit cementite precipitation. As Al is a ferrite stabilizer, which strongly accelerates the ferrite reaction, Mn and Cr contents were increased to avoid the formation of polygonal ferrite during slow air cooling.
The meaning of appearance of the main alloying elements in B1, B2 and B3 steels can be summarized as follows:
C: Intermediate carbon contents were used to depress the bainitic start temperature, with the subsequent improvement in microstructural refinement.
Si/Al: High silicon or aluminum contents were used to inhibit cementite precipitation during austenite decomposition.
Cr and Mo: These elements in combination with Ni were used to increase hardenability. Basically they were intended to avoid the ferrite and perlite formation during air cooling. Other aim was the bainitic start temperature depression to improve microstructural refinement.
Ni: As Cr and Mo, this element was used to increase hardenability. Additionally, it improves toughness when present at high concentrations.
Mn: This element content was kept low as possible to avoid the formation of large blocks of retained austenite.
From calculations performed with the metallurgical model, it was estimated that B1, B2 and B3 steels would present a mainly bainitic microstructure after cooling at 0.1-0.5° C./sec. At the lower end of this range some ferrite was expected to be formed. But even cooling at 0.1° C./sec, the maximum ferrite volume fraction was estimated to be lower than 5% due to the sluggish reaction kinetics associated to the high alloying additions. Conversely, for cooling rates higher than 0.5° C./sec, some martensite was expected to appear. Its maximum amount was difficult to estimate due to uncertainties in the calculation of bainite reaction kinetics. Thereafter, it was expected a final microstructure mainly composed of bainite for cooling rates between 0.1° C./sec and 0.5° C./sec. For these alloys, the calculated bainitic start temperatures (BS) were below 500° C.: 471° C. for B1, 446° C. for B2 and 423° C. for B3. A low transformation temperature was desired to produce an ultrafine structure capable of achieving high strength without loosing toughness.
The bainitic steels B1, B2 and B3 were laboratory melted in a 20 Kg vacuum induction furnace. The obtained steel chemistries are shown in Table 2.
The resulting slabs of 140 mm thickness were hot rolled in a pilot mill to a final thickness of 16 mm. During hot rolling, the reheating and finishing temperatures were 1200-1250° C. and 1000-950° C., respectively. After hot rolling, the plates were air cooled to room temperature.
The as rolled microstructures were analyzed using optical and scanning electron microscopes. Vickers hardness measurements were also performed, and the amount of retained austenite was determined using X-ray diffractometry. Standard tensile and Charpy tests were conducted on as rolled samples. Tensile properties were averaged over results obtained for two samples. Impact properties at room temperature, 0° C. and −20° C. correspond to average values over 3 full-scale Charpy tests for each temperature. In all the cases the samples were taken in the transversal direction.
The continuous cooling transformation diagrams (CCT) of B1, B2 and B3 steels were determined from dilatometric tests performed at a thermomechanical simulator. Cooling rates in the range between 0.1° C./sec and 5° C./sec were considered.
The obtained microstructures were characterized by optical microscopy and hardness measurements.
Several heat treatments were performed on B2 as rolled plates of 16 mm thickness:
From the heat-treated plates tensile and full size Charpy specimens were machined and tested using a procedure similar to that already described for the as rolled materials.
From dilatometric measurements, the CCT diagrams of B1, B2 and B3 steels were derived. In all cases the samples were reheated at 5° C./sec up to 1000° C. without holding time, and then cooled to room temperature at a constant rate (0.1-10° C./sec). For this austenization condition, the austenitic grain size prior to transformation was 40-60 μm for all the materials. The obtained diagrams are presented in
In
From dilatometric data, hardness measurements, optical and scanning electron microscopy, the general phase transformation behavior of B1, B2 and B3 steels was assessed:
B1 steel: For cooling rates higher than 2° C./sec an important part of the transformation took place at temperatures below the calculated martensitic start temperature (MSAndrews=349° C.; see K. W. Andrews, “Empirical formulae for the calculation of some transformation temperatures”, Journal of the Iron and Steel Institute, July 1965, pp. 721).
At 2° C./sec the martensitic volume fraction was estimated from dilatometric measurements as 70%, increasing to 90% at 10° C./sec. These results were supported by optical and scanning electron microscopy. Additionally, the hardness values of the samples cooled at 2, 5 and 10° C./sec were slightly below that corresponding to a fully martensitic structure (
For cooling rates lower than 2° C./sec the transformation start temperature (temperature to 5% of transformation) gradually increased until it reached what it seemed to be a plateau at 500° C. (
For cooling rates between 1.5° C./sec and 0.5° C./sec the final microstructure was mainly bainitic, with retained austenite replacing the M3C carbides.
At cooling rates lower than 0.2° C./sec, some ferritic regions were observed (volume fraction lower than 5%).
B2 steel: It can be seen in
For cooling rates lower than 1° C./sec, the amount of austenite transformed at temperatures above the MS increased continuously. Below 0.5° C./sec great part of the transformation took place above the MS and below the calculated bainitic start temperature (BS=446° C.). For this range of cooling rates the observed microstructure was a fine mixture of bainite and retained austenite.
B3 steel: In this case the microstructure was mostly martensitic at cooling rates higher than 0.8° C./sec. The CCT diagram (
When comparing the results obtained for B1 and B2 steels in
Regarding B3 steel, in comparison to B1 and B2 the hardenability was strongly increased by the Cr and Mn addition. The critical cooling rate to obtain 90% martensite was reduced from 2° C./sec in previous alloys to 0.8° C./sec. Even at 0.1° C./sec about 20% of martensite appeared. On the other hand, the effect of Al as a ferrite stabilizer is evident in the transformation start temperatures: for B3 steel cooled at 0.2-0.1° C./sec the transformation begins at 600° C., whereas for B1 and B2 steels cooled at the same rates the reaction starts 100° C. below the above mentioned temperature.
SEM micrographs of B1 steel in the as rolled condition are shown in
SEM micrographs of B2 steel as rolled are shown in
Regarding B3 steel as rolled, its bainitic structure is finer in comparison to B1 and B2. However, some martensitic regions, which were not present in B1 and B2 steels, appeared in this case. The presence of martensite is not desirable in these materials because it is a brittle phase that impairs toughness. The higher hardenability of B3 steel can be ascribed to the increment in Mn and Cr contents. These additions were intended to compensate the Al acceleration effect on the ferrite reaction kinetics, but it caused the appearance of martensite.
The tensile and impact properties measured for B1, B2 and B3 steels as rolled are shown in the following tables.
When comparing the two high Silicon alloys (B1 and B2), it can be seen that B2 steel presented better tensile and impact properties than B1. This improvement in mechanical properties can be ascribed to the microstructural refinement resulting from the higher carbon addition. In particular, it is interesting to notice that impact property results are in opposition with commonly accepted trends regarding toughness dependence on carbon content, and can be related to the Si presence that is preventing carbide precipitation. When the carbide precipitation is inhibited, an increase in carbon content impairs the ferrite reaction kinetic producing microstructural refinement, with the subsequent increase in strength and toughness. Another important effect is that for the higher carbon steel the appearance of blocky austenitic regions, detrimental to toughness, was reduced probably due to the depletion of the transformation to lower temperatures. Regarding B3 steel, the observed high strength in combination with low toughness can be directly associated to the presence of martensite in the as rolled structure. From Tables 3 and 4, it can be seen that advantageously the best combination of mechanical properties corresponded to B2 steel as rolled: 140 ksi of yield strength and 69 Joules of impact energy at room temperature with a ductile to brittle transition temperature of −20° C. The other two materials did not present 100% ductile fractures in Charpy tests at room temperature.
It is important to notice that from the experimental CCT diagram only minor microstructural differences can be expected when B2 steel is cooled at rates between 0.15° C./sec and 0.30° C./sec, which correspond to air cooling of tubes with wall thickness between 16 mm and 8 mm. Thereafter, nearly the same microstructural and mechanical properties can be obtained with one chemistry for a wide range of tube geometries.
For tubes thicker than 16 mm, and up to 24 mm, the cooling rate at the exit of the hot rolling mill is expected to be in the range between 0.15° C./sec and 0.10° C./sec. In this case some ferrite may be formed. In order to avoid this, an advantageous controlled cooling with cooling rate between 0.2 and 0.5° C./sec can be performed after hot rolling or a chemical composition change.
In summary, B2 steel in the as rolled condition advantageously presented a good combination of tensile and impact properties. In order to further improve strength and toughness, chemical changes or heat treatments are needed.
To study the affect of different microstructural parameters on B2 mechanical properties, several heat treatments were performed, including normalization and tempering at temperatures between 200° C. and 500° C. Some of the results obtained are presented in the following Tables 5 and 6.
In Table 5, it can be seen that the tensile properties were not strongly changed by the normalizing treatment.
The results obtained after tempering were more interesting. There was a strong increase in yield strength after heat treatment at 300° C. The strength improvement can be ascribed to transition carbide precipitation and pinning of dislocations by interstitials. When the tempering temperature was increased to 500° C., the tensile strength decreased (in comparison to the previous treatment) probably due to the replacement of the fine transition carbides by coarse cementite particles. In order to be sure about the metallurgical mechanism that produced the important yield strength increase after tempering at 300° C., a TEM study has been carried out on selected B2 tempered and as rolled samples. A TEM micrograph of B2 steel tempered at 300° C., illustrated in
Regarding impact properties, when comparing the results for the as rolled and normalized materials it is clear that the refinement of the austenitic grain size (from 50-60 μm in the as rolled condition to <30 μm after normalizing) did not produce a toughness improvement. This lack of sensibility to the grain size shows that there is another microstructural parameter (for example the size/thickness of the austenitic regions between the bainitic ferrite laths) that is the toughness controlling factor. The results obtained for the material tempered at 300° C. pointed in the same direction. In this case the toughness was improved without refining the austenitic grains or reducing the bainitic packet size. Considering that the maximum volume fraction of bainite corresponding to the isothermal transformation at 300° C. was not probably achieved during continuous cooling, the retained austenite may continue the reaction during tempering, reducing the size of the interlath austenitic regions. This possibility was supported by:
1) a small reduction of retained austenite, from 13% in the as rolled material to 10% in the sample tempered at 300° C. observed using X-Ray diffractometry; and
2) the progress of the austenite decomposition during tempering at 300° C. observed in a dilatometric test.
Regarding the possibility of a toughness improvement related to the tempering of small martensitic regions, this is not in agreement with the hardness increase observed after tempering at 200-300° C. the as rolled material (see
About the tempering treatment at 500° C., it is clear that extensive carbide precipitation and coarsening took place, slightly improving yield strength but strongly deteriorating toughness in comparison to the as rolled material.
The results obtained after different heat treatments of B2 steel showed that one way to improve yield strength and toughness is by tempering the material at low temperatures comprised between 200 and 350° C., preferably about 300° C. In this case, the transition carbide precipitation improves strength, and the refinement in the bainitic microstructure and the reduction in the interlath austenitic regions improve toughness.
In
The most promising combination of mechanical properties was obtained with B2 steel tempered at 300° C. Due to the high yield strength and good toughness, this material is positioned above the toughness-strength curve of quenched and tempered steels.
From the results obtained, it can be concluded that the bainitic steel of the invention in the as rolled condition has good combination of strength and toughness when the microstructure is composed of a fine mixture of bainitic ferrite and retained austenite (B2 steel). If the structure is coarse with blocks of retained austenite between bainitic sheaves (B1 steel) or when large martensitic regions are present (B3 steel) the impact properties are impaired.
With its fine bainitic structure without large blocky austenitic or martensitic regions B2 steel as rolled is therefore suitable for OCTG applications.
The most promising combination of mechanical properties was obtained with B2 steel as rolled and tempered at 300° C. Due to the high yield strength and good toughness, this material is positioned above the toughness-strength curve of the quenched and tempered martensitic steels.
Advantageously the bainitic steel tubes or pipes, obtained by means of the process of the invention, have homogeneous mechanical properties due to the avoidance of the quenching treatment. In particular B2 steel, hot rolled and tempered, presents the same mechanical properties for a wide range of tube wall thickness, between 18 mm and 8 mm.
For tubes up to 18 mm thickness, the alloying additions in B2 steel can be reduced if accelerated cooling after hot rolling is available.
For thicker tubes (up to 35 mm), the decrease in the cooling rate at the exit of the hot rolling mill has to be compensated by a controlled cooling at 0.10-1.0° C./sec, preferably 0.2-0.5° C.; or by alloying additions.
Modifications of B2 steel chemistry may be performed without changing the principles of the invention, that is to produce an ultra-fine bainitic structure in the as rolled condition with minor fractions of martensite and blocky austenitic regions, and, in a advantageous embodiment of the invention, to perform a tempering at low temperature to increase the yield to tensile strength ratio to make the material suitable for high strength OCTG applications. For example Ni can be substituted by Mn as an austenitizing element, Cr and C contents may be changed depending on tube thickness, or microalloying elements (Ti and Nb) may be added to control austenitic grain size during hot rolling.
It is to be understood that the foregoing description and specific embodiments are merely illustrative of the best mode of the invention and the principles thereof, and that various modifications and additions may be made by those skilled in the art, without departing from the spirit and scope of this invention, which is therefore understood to be limited only by the scope of the appended claims.
| Filing Document | Filing Date | Country | Kind | 371c Date |
|---|---|---|---|---|
| PCT/EP2007/062492 | 11/19/2007 | WO | 00 | 8/9/2010 |
