Field of the Invention

This invention relates to high strength steel and its manu¬ facture, the steel being useful in structural applications as well as being a precursor for linepipe. More particularly, this invention relates to the manufacture of dual phase, high strength steel plate comprising ferrite and martenβite/bainite phases wherein the micro- structure and mechanical properties are substantially uniform through the thickness of the plate, and the plate is characterized by superior toughness and weldability.

Background of the Invention

Dual phase steel comprising ferrite, a relatively soft phase and martenβite/bainite, a relatively strong phase, are produced by annealing at temperatures between the A _r 3 and A _r - transformation points, followed by cooling to room temperature at rates ranging from air cooling to water quenching. The selected annealing temperature is dependent on the the steel chemistry and the desired volume relation¬ ship between the ferrite and martensite/bainite phases.

The development of low carbon and low alloy dual phase steels is well documented and has been the subject of extensive research in the metallurgical community; for example, conference proceedings on "Fundamentals of Dual Phase Steels" and "Formable HSLA and Dual Phase Steels", U.S. patents 4,067,756 and 5,061,325. However, the applica¬ tions for dual phase steels have been largely focused on the auto¬ motive industry wherein the unique high work hardening characteristics of this steel are utilized for promoting formability of automotive sheet steels during processing and stamping operations. Consequently, dual phase steels have been limited to thin sheets, typically in the range of 2-3 mm, and less than 10 mm, and exhibit yield and ultimate tensile strengths in the range of 50-60 ksi and 70-90 ksi, respec¬ tively. Also, the volume of the martensite/bainite phase generally

represents about 10-40% of the microstructure, the remainder being the softer ferrite phase.

Consequently, an object of this invention is utilizing the high work hardening capability of dual phase steel not for improving formability, but for achieving rather high yield strengths, after the 1-3% deformation imparted to plate steel during the formation of linepipe to > 100 ksi, preferably > 110 ksi. Thus, dual phase steel plate having the characteristics to be described herein is a precursor for linepipe.

An object of this invention is to provide substantially uniform microstructure through the thickness of the plate for plate thickness of at least 10 mm. A further object is to provide for a fine scale distribution of constituent phases in the microstructure so as to expand the useful boundaries of volume percent bainite/ martensite to about 75% and higher, thereby providing high strength, dual phase steel characterized by superior toughness. A still further object of this invention is to provide a high strength, dual phase steel having superior weldability and superior heat affected zone (HAZ) softening resistance.

Summary of the Invention

In accordance with this invention , steel chemistry is balanced with thermomechanical control of the rolling process, thereby allowing the manufacture of high strength, i.e., yield strengths greater than 100 ksi, and at least 110 ksi after 1-3% deformation, dual phase steel useful as a precursor for linepipe, and having a microstructure comprising 40-80%, preferably 50-β0% by volume of a martensite/bainite phase in a ferrite matrix, the bainite being less than about 50% of martensite/bainite phase.

In a preferred embodiment, the ferrite matrix is further strengthened with a high density of dislocations, i.e., >10 ¹⁰ cm/cm ³ , and a dispersion of fine sized precipitates of at least one and preferably all of vanadium and niobium carbides or carbonitrides, and

molybdenum carbide, i.e., (V,Nb)(C,N) and M02C. The very fine (< 5θA diameter) precipitates of vanaditim, niobium and molybdenum carbides or carbonitrides are formed in the ferrite phase by interphase precipita¬ tion reactions which occur during austenite ferrite transformation below the A--3 temperature. The precipitates are primarily vanadium and niobium carbides and are referred to as (V,Nb)(C,N). Thus, by balancing the chemistry and the thermomechanical control of the rolling process, dual phase steel can be produced in thicknesses of at least about 15 mm, preferably at least about 20 mm and having ultr- ahigh strength.

The strength of the steel is related to the presence of the martenβite/bainite phase, where increasing phase volume results in increasing strength. Nevertheless, a balance must be maintained between strength and toughness (ductility) where the toughness is provided by the ferrite phase. For example, yield strengths after 2% deformation of at least about 100 kβi are produced when the martensite/bainite phase is present in at least about 40 vol%, and at least about 120 ksi when the martensite/bainite phase is at least about 60 vol%.

The preferred steel, that is, with the high density of dislocations and vanadium and niobium precipitates in the ferrite phase is produced by a finish rolling reduction at temperatures between the A _r 3 and A _r χ transformation points and quenching to room temperature. The procedure, therefore, is contrary to dual phase steels for the automotive industry, usually 10 mm or less thickness and 50-60 ksi yield strength, where the ferrite phase must be free of precipitates to ensure adequate formability. The precipitates form discontinuously at the moving interface between the ferrite and austenite. However, the precipitates form only if adequate amounts of vanadium or niobium or both are present and the rolling and heat treatment conditions are carefully controlled. Thus, vanadium and niobium are key elements of the steel chemistry.

DESCRIPTION OF THE DRAWINGS

Figure 1 shows a scanning electron micrograph revealing ferrite phase (grey) and martensite/bainite phase (brighter region) alloy A3 quench. Thiβ figure shows the final product of the dual phase steel produced in accordance with this invention.

Figure 2 shows a transmissions electron micrograph of niobium and vanadium carbonitride precipitates in the range of less than about 50λ, preferably about lθ-5θA, in the ferrite phase.

Figures 3a and 3b show transmission electron micrographs of the microstructural detail of the strong phase martensite. Figure 3a is a bright field image, and Figure 3b a dark field image correspond¬ ing to Figure 3a.

Figure 4 shows plots of hardness (Vickers) data across the HAZ (ordinate) for the steel produced by this invention (solid line) and a similar plot for a commercial X100 linepipe steel (dotted line). The steel of this invention shows no significant decrease in the HAZ strength, whereas a significant decrease, approximately 15%, in HAZ strength (as indicated by the Vickers hardness) occurs for the X100 steel.

Now, the steel of this invention provides high strength superior weldability and low temperature toughness and comprises, by weight:

0.05 - 0.12% C, preferably 0.06 - 0.12, more preferably 0.07 - 0.09

0.01 - 0.5% Si

0.4 - 2.0% Mn, preferably 1.0 - 2.0, more preferably 1.2 - 2.0

0.03 - 0.12% Nb, preferably 0.05 - 0.1

0.05 - 0.15% V

0.2 - 0.8% Mo

0.3 - 1.0% Cr, preferred for hydrogen containing environments

0.015- 0.03% Ti

0.01 - 0.03% Al

Pcm < 0.24

the balance being Fe and incidental impuritieβ.

The sum of the vanadium and niobium concentrations is > 0.1 wt%, and more preferably vanadium and niobium concentrations each are > 0.04%. The well known contaminants N, P, S are minimized even though some N is desired, as explained below, for producing grain growth inhibiting titanium nitride particles. Preferably, N concen¬ tration is about 0.001-0.01 wt%, S no more than 0.01 wt%, and P no more than 0.01 wt%. In this chemistry the steel is boron free in that there is no added boron, and boron concentration is < 5 ppm, prefer¬ ably < 1 ppm.

Generally, the material of thiβ invention is prepared by forming a steel billet of the above composition in normal fashion; heating the billet to a temperature sufficient to dissolve substan¬ tially all, and preferably all vanadium carbonitrides and niobium carbonitrides, preferably in the range of 1150-1250°C. Thus essen¬ tially all of the niobium, vanadium and molybdenum will be in solu¬ tion; hot rolling the billet in one or more passes in a first reduction providing about 30-70% reduction at a first temperature range where austenite recrystallizes; hot rolling the reduced billet in one or more passes in a second rolling reduction providing about 40-70% reduction in a second and somewhat lower temperature range when austenite does not recrystallize but above the Ar3; air cooling to a temperature in the range between A _r 3 and A _r ι transformation points and where 20-60% of the austenite haβ tranβformed to ferrite; rolling the further reduced billet in one or more passes in a third rolling reduction of about 15-25%; water cooling at a rate of at least 25°C/βecond, preferably at least about 35°C/βecond, thereby hardening the billet, to a temperature no higher than 400°C, where no further transformation to ferrite can occur and, if desired, air cooling the rolled, high strength steel plate, useful as a precursor for linepipe to room temperature. As a result, grain size is quite uniform and < 10 microns, preferably < 5 microns.

High strength steels necessarily require a variety of proper¬ ties and these properties are produced by a combination of elements and mechanical treatments. The role of the various alloying elements and the preferred limits on their concentrations for the present invention are given below:

Carbon provides matrix strengthening in all steels and welds, whatever the microstructure, and also precipitation strengthening through the formation of small NbC and VC particles, if they are sufficiently fine and numerous. In addition, NbC precipitation during hot rolling serves to retard recrystallization and to inhibit grain growth, thereby providing a means of austenite grain refinement. This leads to an improvement in both strength and low temperature tough¬ ness. Carbon also assists hardenability, i.e., the ability to form harder and stronger microstructures on cooling the steel. If the carbon content is less than 0.01%, these strengthening effects will not be obtained. If the carbon content is greater than 0.12%, the steel will be susceptible to cold cracking on field welding and the toughness is lowered in the steel plate and its heat affected zone (HAZ) on welding.

Manganese is a matrix βtrengthener in steels and welds and it also contributes strongly to the hardenability. A minimum amount of 0.4% Mn is needed to achieve the necessary high strength. Like carbon, it is harmful to toughness of plates and welds when too high, and it also causes cold cracking on field welding, so an upper limit of 2.0% Mn is imposed. This limit is also needed to prevent severe center line segregation in continuously cast linepipe steels, which is a factor helping to cause hydrogen induced cracking (HIC) .

Silicon is always added to steel for deoxidization purposes and at least 0.01% is needed in this role. In greater amounts Si has an adverse effect on HAZ toughness, which is reduced to unacceptable levels when more than 0.5% is present.

Niobium is added to promote grain refinement of the rolled microstructure of the steel, which improves both the strength and the

toughness. Niobium carbide precipitation during hot rolling serves to retard recryβtallization and to inhibit grain growth, thereby provid¬ ing a meanβ of austenite grain refinement. It will give additional strengthening on tempering through the formation of NbC precipitates. However, too much niobium will be harmful to the weldability and HAZ toughness, so a maximum of 0.12% is imposed.

Titanium, when added as a small amount is effective in forming fine particles on TiN which refine the grain size in both the rolled structure and the HAZ of the steel. Thus, the toughness is improved. Titanium is added in such an amount that the ratio Ti/N ranges between 2.0 and 3.4. Excess titanium will deteriorate the toughness of the steel and welds by forming coarser TiN or Tie particles. A titanium content below 0.002% cannot provide a suffi¬ ciently fine grain size, while more than 0.04% causes a deterioration in toughness.

Aluminum is added to these steels for the purpose of de- oxidization. At least 0.002% Al is required for this purpose. If the aluminum content is too high, i.e., above 0.05%, there is a tendency to form AI2O3 type inclusions, which are harmful for the toughness of the steel and its HAZ.

Vanadium is added to give precipitation strengthening, by forming fine VC particles in the steel on tempering and its HAZ on cooling after welding. When in solution, vanadium is potent in promoting hardenability of the steel. Thus vanadium will be effective in maintaining the HAZ strength in a high strength steel. There is a maximum limit of 0.15% since excessive vanadium will help cause cold cracking on field welding, and also deteriorate the toughness of the steel and its HAZ. Vanadium is also a potent βtrengthener to eutectoidal ferrite via interphase precipitation of vanadium carbo¬ nitride particleβ of < about 5θA diameter, preferably 10-50A diameter.

Molybdenum increases the hardenability of a steel on direct quenching, so that a strong matrix microstructure is produced and it also gives precipitation strengthening on reheating by forming M02C

and NbMo particles. Excessive molybdenum helps to cause cold cracking on field welding, and also deteriorate the toughness of the steel and HAZ, so a maximum of 0.8% is specified.

Chromium also increases the hardenability on direct quench¬ ing. It improves corrosion and HIC resistance. In particular, it is preferred for preventing hydrogen ingress by forming a Cr2θ3 rich oxide film on the steel surface. As for molybdenum, excessive chromium helps to cause cold cracking on field welding, and also deteriorate the toughness of the steel and its HAZ, so a maximum of 1.0% Cr is imposed.

Nitrogen cannot be prevented from entering and remaining in steel during βteelmaking. In thiβ steel a small amount is beneficial in forming fine TiN particles which prevent grain growth during hot rolling and thereby promote grain refinement in the rolled steel and its HAZ. At least 0.001% N is required to provide the necessary volume fraction of TiN. However, too much nitrogen deteriorates the toughness of the steel and its HAZ, so a maximum amount of 0.01% N is imposed.

The objectives of the thermomechanical processing are two fold: producing a refined and flattened austenitic grain and intro¬ ducing a high density of dislocations and shear bands in the two phases.

The first objective is satisfied by heavy rolling at tempera¬ tures above and below the austenite recrystallization temperature but always above the A--3. Rolling above the recrystallization temperature continuously refines the austenite grain size while rolling below the recrystallization temperature flattens the austenitic grain. Thus, cooling below the A _r 3 where austenite begins its transformation to ferrite results in the formation of a finely divided mixture of austenite and ferrite and, upon rapid cooling below the A _r ^, to a finely divided mixture of ferrite and martensite/bainite.

The second objective is satisfied by the third rolling reduction of the flattened austenite grains at temperatures between the A _r ι and A _r 3 where 20% to 60% of the austenite has transformed to ferrite.

The thermomechanical processing practiced in this invention is important for inducing the desired fine distribution of constituent phases.

The temperature that defines the boundary between the ranges where austentite recrystallizes and where austenite does not re- crystallize depends on the heating temperature before rolling, the carbon concentration, the niobium concentration and the amount of reduction in the rolling passes. This temperature can be readily determined for each steel composition either by experiment or by model calculation.

Linepipe is formed from plate by the well known U-O-E process in which plate is formed into a U shape, then formed into an O shape, and the O shape is expanded 1-3%. The forming and expansion with their concommitant work hardening effects leads to the highest strength for the linepipe.

The following examples illustrate the invention described herein.

A 500 lb. heat of the alloy represented by the following chemistry waβ vacuum induction melted, caβt into ingots, forged into 4 inch thick slabs, heated at 1240°C for two hours and hot rolled according to the schedule in Table 2.

TABLE 1

Chemical Composition ιwt%)

S Mn Si Mo Cr Nb

0.074 1.58 0.13 0.30 0.34 0.086

V Jl. S 2. Nfppffi) ^p cm

0.082 0.020 0.026 0.006 0.006 52 0.20

The alloy and the thermomechanical procesβing were designed to produce the following balance with regard to the strong carbonitride formers, particularly niobium and vanadium:

• about one third of these compounds precipitate in austenite prior to quenching; these precipitates provide recrystallization resis¬ tance as well as austenite grain pinning resulting in fine austenite grains before it transforms;

• about one third of these compounds precipitate during austenite to ferrite transformation through the intercritical and subcritical region; these precipitates help strengthen the ferrite phase;

• about one third of these compounds are retained in solid solution for precipitation in the HAZ and ameliorateing or eliminating the normal softening seen with other steels.

The thermomechanical rolling schedule for the 100 mm square initial forged slab is shown below:

TABLE 2

Starting Thickness: 100 ]mm

Reheat Temperature: 1240 ^β C

Reheating Time: 2 hours

Thickness After Temperature

Pass Pass, mm °C

0 100 1240

1 85 1104

2 70 1082

3 57 1060

Delay (turn piece on edge) (1)

4 47 899

5 38 866

6 32 852

7 25 829

»_ I". 1 atj> /•r-nr'n ι- *J Ar*A aHni

Immediately Water Quench To Room Temperature (2)

(1) Delay amounted to air cooling, typically at about 1°C/second.

(2) Quenching rate from finish temperature should be in the range 20 to 100°C/second and more preferably, in the range 30 to 40°C/second to induce the desired dual phase microstructure in thick sections exceeding 20 mm in thickness.

The final product waβ 20 mm thick and waβ 45% ferrite and 55% martenβite/bainite.

To vary the amounts of ferrite and the other austenite decompoβition productβ, quenching from variouβ finiβh temperatures was conducted as described in Table 3. The ferrite phase includes both the proeutectoidal (or "retained ferrite") and the eutectoidal (or "transformed" ferrite) and signifies the total ferrite volume fraction. When the steel waβ quenched from 800°C, it waβ in the 100% austenite region, indicating that the Ar3 temperature is below 800°C. Aβ βeen from Figure 1, the austenite is 75% transformed when quenching

from about 725°C, indicating that the r^ temperature is close to this temperature, thus indicating a two phase window for this alloy of about 75°C. Table 3 summarizes the finish rolling, quenching, volume fractions and the Vickers microhardneββ data.

TABES 3 Dual Phase Microstructures and TMCP Practice

Finish Start %

Alloy Roll Quench % Martensite/ Hardness .1) Temp .°C) Temp ι°C) Ferrite Bainite fHV)

Al 800 800 0 100 260

A2 750 750 45 55 261

A3 750 740 60 40 261

A4 725 725 75 25 237

(1) composition shown in Table 1.

Because steels having a high volume percentage of the second or martenβite/bainite phase are usually characterized by poor ductility and toughness, the steels of this invention are remarkable in maintaining sufficient ductility to allow forming and expansion in the UOE process. Ductility is retained by maintaining the effective dimensions of microstructural units such as the martensite packet below 10 microns and the individual features within this packet below 1 micron. Figure 1, the scanning electron microscope (SEM) micro¬ graph, shows the dual phase microstructure containing ferrite and martensite for processing condition A3. Remarkable uniformity of microstructure throughout the thickness of the plate was observed in all dual phase steels.

Figure 2 shows a transmission electron micrograph revealing a very fine dispersion of interphase precipitates in the ferrite region of A3 steel. The eutectoidal ferrite is generally observed close to the interface at the second phase, dispersed uniformly throughout the sample and its volume fraction increaβeβ with lowering of the tempera¬ ture from which the steel is quenched.

Figures 3a and 3b show transmission electron micrographs revealing the nature of the second phase in these steels. A

predominantly lath martenβitic microstructure with some bainitic phase was observed. The martensite revealed thin film, i.e., less than about 500 A thick, retained austenite at the lath boundaries as shown in the dark field image, Figure 3b. This morphology of martensite ensures a βtrong but alβo a tough second phase contributing not only to the strength of the two phase steel but also helping to provide good toughness.

Table 4 shows the tensile strength and ductility of two of the alloy A samples.

Yield strength after 2% elongation in pipe forming will meet the minimum deβired strength of at least 100 kβi, preferably at least 110 kβi, due to the excellent work hardening characteriβticβ of theβe microstructures.

Table 5 showβ the Charpy-V-Notch impact toughness (ASTM specification E-23) at -40 and -76°C performed on longitudinal (L-T) samples of alloy A4.

TABLE 5

% Ferrite/ Test Temperature Alloy % Martensite CO Energy (Joules)

A4 75/25 -40 301

-76 269

The impact energy values captured in the above table indicate excellent toughness for the steels of this invention. The βteel of this invention haβ a toughness of at least 100 joules at -40 ^β C, preferably at least about 120 joules at -40°C.

A key aspect of the present invention is a high strength steel with good weldability and one that has excellent HAZ softening resistance. Laboratory single bead weld tests were performed to observe the cold cracking susceptibility and the HAZ softening. Figure 4 presents an example of the data for the steel of this inven¬ tion. This plot dramatically illustrates that in contrast to the βteelβ of the βtate of the art, for example commercial X100 linepipe βteel, the dual phaβe βteel of the preβent invention, doeβ not βuffer from any significant or measurable softening in the HAZ. In contrast X100 showβ a 15% softening as compared to the base metal. By follow¬ ing this invention the HAZ haβ at least about 95% of the strength of the base metal, preferably at least about 98% of the strength of the base metal. These strengths are obtained when the welding heat input ranges from about 1-5 kilo joules/mm.