Patent application title: HIGH STRENGTH THIN CAST STRIP PRODUCT AND METHOD FOR MAKING THE SAME
Kristin R. Carpenter (Fairy Meadow, AU)
Scott Cluff (Fairy Meadow, AU)
Christopher R. Killmore (Wollongong, AU)
Daniel Geoffrey Edelman (Brownsburg, IN, US)
Daniel Geoffrey Edelman (Brownsburg, IN, US)
Peter C. Campbell (Zionsville, IN, US)
Walter N. Blejde (Brownsburg, IN, US)
IPC8 Class: AB22D1106FI
Class name: With casting or solidifying from melt iron(fe) or iron base alloy continuous casting
Publication date: 2014-01-16
Patent application number: 20140014238
A steel product or thin steel cast strip including, by weight, less than
0.25% carbon, between 0.20 and 2.0% manganese, between 0.05 and 0.50%
silicon, less than 0.01% aluminum, niobium between 0.01% and 0.20%, and
vanadium between 0.01% and 0.20%. The steel product may have a tensile
strength of between 550 and 600 MPa after age hardening at a peak
temperature of between 700 and 750° C.
1. An age hardened steel product comprising, by weight, less than 0.25%
carbon, between 0.20 and 2.0% manganese, between 0.05 and 0.50% silicon,
less than 0.01% aluminum, at least one element from the group consisting
of niobium between about 0.01% and about 0.20%, vanadium between about
0.01% and about 0.20%, and having a tensile strength of at least 550 MPa
after age hardening at a peak temperature of between 700 and 750.degree.
2. The age hardened steel product of claim 1 where the tensile strength is between 550 and 600 MPa.
3. A method of preparing a thin cast steel strip comprising the steps of: (a) assembling an internally cooled roll caster having laterally positioned casting rolls forming a nip between them, and forming a casting pool of molten steel supported on the casting rolls above the nip and confined adjacent the ends of the casting rolls by side dams, (b) counter rotating the casting rolls to solidify metal shells on the casting rolls as the casting rolls move through the casting pool, (c) forming steel strip from the metal shells cast downwardly through the nip between the casting rolls having a composition comprising by weight, less than 0.25% carbon, less than 0.01% aluminum, and at least one element from the group consisting of niobium between about 0.01% and about 0.20%, vanadium between about 0.01% and about 0.20%, and (d) age hardening the steel strip at a peak temperature between 700.degree. C. and 750.degree. C. for a duration of between 9 and 11 seconds to form an age hardened thin cast steel strip having a tensile strength of at least 550 MPa.
4. The method of claim 3 where the tensile strength of the age hardened thin cast steel strip is between 550 and 600 MPa.
 This is a nonprovisional application claiming benefit from U.S. Provisional Application No. 61/672,042, filed on Jul. 16, 2012, which is incorporated herein by reference.
 This invention relates to making of high strength thin cast strip, and the method for making such cast strip by a twin roll caster.
 In a twin roll caster, molten metal is introduced between a pair of counter-rotated, internally cooled casting rolls so that metal shells solidify on the moving roll surfaces, and are brought together at the nip between them to produce a solidified strip product, delivered downwardly from the nip between the casting rolls. The term "nip" is used herein to refer to the general region at which the casting rolls are closest together. The molten metal is poured from a ladle through a metal delivery system comprised of a tundish and a core nozzle located above the nip to form a casting pool of molten metal, supported on the casting surfaces of the rolls above the nip and extending along the length of the nip. This casting pool is usually confined between refractory side plates or dams held in sliding engagement with the end surfaces of the rolls so as to dam the two ends of the casting pool against outflow.
 In the past, high-strength low-carbon thin strip with yield strengths of 413 MPa (60 ksi) or higher, in strip thicknesses less than 3.0 mm, have been made by recovery annealing of cold rolled strip. Cold rolling was required to produce the desired thickness. The cold roll strip was then recovery annealed to improve the ductility without significantly reducing the strength. However, the final ductility of the resulting strip was still relatively low and the strip would not achieve total elongation levels over 6%, which is required for structural steels by some building codes for structural components. Such recovery annealed cold rolled, low-carbon steel was generally suitable only for simple forming operations, e.g., roll forming and bending. Producing this steel strip with higher ductility was not technically feasible in these final strip thicknesses using the cold rolled and recovery annealed manufacturing route.
 In the past, such high strength steel had been made by microalloying with elements such as niobium (Nb), vanadium (V), titanium (Ti) or molybdenum (Mo), and hot rolling to achieve the desired thickness and strength level. Such microalloying required expensive and high levels of niobium, vanadium, titanium or molybdenum and resulted in formation of a bainite-ferrite microstructure typically with 10% to 20% bainite. See, U.S. Pat. No. 6,488,790. Alternatively, the microstructure could be ferrite with 10% to 20% pearlite. Hot rolling the strip resulted in the partial precipitation of these alloying elements. As a result, relatively high alloying levels of the Nb, V, Ti or Mo elements were required to provide enough age hardening of the predominately ferritic transformed microstructure to achieve the required strength levels. These high microalloying levels significantly raised the hot rolling loads needed and restricted the thickness range of the hot rolled strip that could be economically and practically produced. Such alloyed high strength strip could be directly used for galvanizing after pickling for the thicker end of the product range greater than 3 mm in thickness.
 However, making high strength, steel strip less than 3 mm in thickness with additions of Nb, V, Ti or Mo to the base steel chemistry was very difficult, particularly for wide strip due to the high rolling loads, and not always commercially feasible. In the past, large additions of these elements were needed for strengthening the steel, and in addition, caused reductions in elongation properties of the steel. High strength microalloyed hot rolled strips in the past were relatively inefficient in providing strength, relatively expensive, and often required compensating additions of other alloying elements.
 Additionally, cold rolling was generally required for lower thicknesses of strip; however, the high strength of the hot rolled strip made such cold rolling difficult because of the high cold roll loadings required to reduce the thickness of the strip. These high alloying levels also considerably raised the recrystallization annealing temperature needed, requiring expensive to build and to operate annealing lines capable of achieving the high annealing temperature needed for full recrystallization annealing of the cold rolled strip.
 In short, the application of previously known microalloying practices with Nb, V, Ti or Mo elements to produce high strength thin strip could not be commercially economically produced because of the high alloying costs, relative inefficiency of element additions, difficulties with high rolling loads in hot rolling and cold rolling, and the high recrystallization annealing temperatures required.
 Presently disclosed is an age hardened steel product comprising, by weight, less than 0.25% carbon, between 0.20 and 2.0% manganese, between 0.05 and 0.50% silicon, less than 0.01% aluminum, at least one element from the group consisting of niobium between about 0.01% and about 0.20%, vanadium between about 0.01% and about 0.20%, and having a tensile strength of at least 550 MPa after age hardening at a peak temperature of between 700 and 750° C.
 The age hardened steel product may have a tensile strength between 550 and 600 MPa.
 Also disclosed is a method of preparing a thin cast steel strip comprising the steps of:
 (a) assembling an internally cooled roll caster having laterally positioned casting rolls forming a nip between them, and forming a casting pool of molten steel supported on the casting rolls above the nip and confined adjacent the ends of the casting rolls by side dams,
 (b) counter rotating the casting rolls to solidify metal shells on the casting rolls as the casting rolls move through the casting pool,
 (c) forming steel strip from the metal shells cast downwardly through the nip between the casting rolls having a composition comprising by weight, less than 0.25% carbon, less than 0.01% aluminum, and at least one element from the group consisting of niobium between about 0.01% and about 0.20%, vanadium between about 0.01% and about 0.20%, and
 (d) age hardening the steel strip at a peak temperature between 700° C. and 750° C. for a duration of between 9 and 11 seconds to form an age hardened thin cast steel strip having a tensile strength of at least 550 MPa.
 The tensile strength of the age hardened thin cast steel strip may be between 550 and 600 MPa.
BRIEF DESCRIPTION OF THE DRAWINGS
 FIG. 1 illustrates a strip casting installation incorporating an in-line hot rolling mill and coiler;
 FIG. 2 illustrates details of the twin roll strip caster;
 FIG. 3 is a graph showing the continuous cooling transformation diagram for the plain C--Mn ultra-thin cast strip steel;
 FIG. 4 is a graph showing the continuous cooling transformation diagram for the V ultra-thin cast strip steel;
 FIG. 5 is a graph showing the continuous cooling transformation diagram for the Nb ultra-thin cast strip steel;
 FIG. 6 is a graph showing the continuous cooling transformation diagram for the Nb--V ultra-thin cast strip steel;
 FIGS. 7(a)-(d) are micrographs of the microstructures of the ultra-thin cast strip steels cooled at 30° C./second for the (a) C--Mn, (b) V, (c) Nb, and (d) Nb--V grades;
 FIG. 8 is a graph showing the Vickers hardness results for the C--Mn, V, Nb, and Nb--V grades;
 FIG. 9 is a graph showing the effect of niobium amount on the yield and tensile strengths at varying levels of Mn content;
 FIG. 10 is a graph showing the effect of hot rolling reduction percentage for the 0.043% V and plain carbon base steels;
 FIGS. 11(a)-(b) are micrographs showing the microstructure for (a) typical C--Mn UCS in the grade 380 MPa condition and (b) typical 0.04% V UCS grade;
 FIGS. 12(a)-(d) are micrographs showing the effect of increasing Nb additions on the microstructure of UCS, where (a) shows 0.015% Nb, (b) shows 0.025% Nb, (c) shows 0.04% Nb and (d) shows 0.08% Nb;
 FIG. 13(a)-(b) are photomicrographs showing unrecrystallized autenite grains in the Nb steel after (a) 15% hot rolling reduction and (b) 35% hot rolling reduction;
 FIG. 14 is a graph showing strength as a function of peak aging temperature; and
 FIG. 15 is a graph showing temperature for peak strength as a function of time at peak temperature.
DETAILED DESCRIPTION OF THE DRAWINGS
 FIG. 1 illustrates successive parts of strip caster for continuously casting steel strip. FIGS. 1 and 2 illustrate a twin roll caster 11 that continuously produces a cast steel strip 12, which passes in a transit path 10 across a guide table 13 to a pinch roll stand 14 having pinch rolls 14A. Immediately after exiting the pinch roll stand 14, the strip passes into a hot rolling mill 16 having a pair of reduction rolls 16A and backing rolls 16B where the cast strip is hot rolled to reduce a desired thickness. The hot rolled strip passes onto a run-out table 17 where the strip may be cooled by convection and contact with water supplied via water jets 18 (or other suitable means) and by radiation. The rolled and cooled strip is then passes through a pinch roll stand 20 comprising a pair of pinch rolls 20A and then to a coiler 19. Final cooling of the cast strip takes place after coiling.
 As shown in FIG. 2, twin roll caster 11 comprises a main machine frame 21, which supports a pair of laterally positioned casting rolls 22 having casting surfaces 22A. Molten metal is supplied during a casting operation from a ladle (not shown) to a tundish 23, through a refractory shroud 24 to a distributor or moveable tundish 25, and then from the distributor 25 through a metal delivery nozzle 26 between the casting rolls 22 above the nip 27. The molten metal delivered between the casting rolls 22 forms a casting pool 30 above the nip. The casting pool 30 is restrained at the ends of the casting rolls by a pair of side closure dams or plates 28, which are pushed against the ends of the casting rolls by a pair of thrusters (not shown) including hydraulic cylinder units (not shown) connected to the side plate holders. The upper surface of casting pool 30 (generally referred to as the "meniscus" level) usually rises above the lower end of the delivery nozzle so that the lower end of the delivery nozzle is immersed within the casting pool 30. Casting rolls 22 are internally water cooled so that shells solidify on the moving roller surfaces as they pass through the casting pool, and are brought together at the nip 27 between them to produce the cast strip 12, which is delivered downwardly from the nip between the casting rolls.
 The twin roll caster may be of the kind that is illustrated and described in some detail in U.S. Pat. Nos. 5,184,668 and 5,277,243 or U.S. Pat. No. 5,488,988,or U.S. patent application Ser. No. 12/050,987. Reference may be made to those patents for appropriate construction details of a twin roll caster appropriate for use in an embodiment of the present invention.
 The in-line hot rolling mill 16 is typically used for reductions of 10% to 50%. On the run-out-table 17 the cooling may include water cooling section and air mist cooling to control cooling rates of austenite transformation to achieve desired microstructure and material properties.
 Hot reductions larger than about 20% can induce the recrystallization of austenite, which reduces the grain size and volume fraction of acicular ferrite. We have found that the addition of alloying elements increasing the hardenability of the steel suppressed the recrystallization of the coarse as-cast austenite grain size during the hot rolling process, and resulted in the hardenability of the steel being retained after hot rolling, enabling thinner material to be produced with the desired microstructure and mechanical properties.
 Microalloying elements in steel are commonly taken to refer to the elements titanium niobium, and vanadium. These elements were usually added in the past in levels below 0.1%, but in some cases in levels as high as 0.2%. These elements are capable of exerting strong effects on the steel microstructure and properties via a combination of hardenability, grain refining and strengthening effects (in the past as carbonitride formers). Molybdenum has not normally been regarded as a microalloying element since on its own it is a relatively weak carbonitride former, but may be effective in the present circumstances and may form complex carbonitride particles along with niobium and vanadium. Carbonitride formation is inhibited in the hot rolled strip with these elements as explained below.
 The high strength thin cast strip product combines several attributes to achieve a high strength light gauge cast strip product by microalloying with these elements. Strip thicknesses may be less than 3 mm, less than 2.5 mm, or less than 2.0 mm, and may be in a range of 0.5 mm to 2.0 mm. The cast strip is produced by hot rolling without the need for cold rolling to further reduce the strip to the desired thickness. Thus, the high strength thin cast strip product overlaps both the light gauge hot rolled thickness ranges and the cold rolled thickness ranges desired. The strip may be cooled at a rate of 10° C. per second and above, and still form a microstructure that is a majority and typically predominantly bainite and acicular ferrite.
 The benefits achieved through the preparation of such a high strength thin cast strip product are in contrast to the production of previous conventionally produced microalloyed steels that result in relatively high alloy costs, inefficiencies in microalloying, difficulties in hot and cold rolling, and difficulties in recrystallization annealing since conventional continuous galvanizing and annealing lines are not capable of providing the high annealing temperatures needed. Moreover, the relatively poor ductility exhibited with strip made by the cold rolled and recovery annealed manufacturing route is overcome.
 In previous conventionally produced microalloyed steels, elements such as niobium and vanadium could not remain in solid solution through solidification, hot rolling, coiling and cooling. The niobium and vanadium diffused through the microstructure forming carbonitride particles at various stages of the hot coil manufacturing process. Carbonitride particles, in the present specification and appended claims, includes carbides, nitrides, carbonitrides, and combinations thereof. The formation and growth of carbon and nitrogen particles in the hot slab and subsequent coiling of previous conventionally produced microalloyed steels further reduced the grain size of austenite in the hot slab, decreasing the hardenability of the steel. In these prior steels, the effect of particles in the hot slab had to be overcome by increasing the amount of microalloying elements, reheating the cast slabs to higher temperatures, and lowering carbon content.
 The steels studied for the CCT diagrams are listed in Table I, where the base or reference grade is a plain C--Mn steel. The C, Mn and Si levels were similar for all four steels, while the Nb and V levels in the microalloyed steels were about 0.04% to allow direct comparison of the CCT diagrams between the Nb, V and Nb+V microalloyed steels. The construction of a CCT diagram is a technique to estimate the influence of alloy additions on the austenite transformation and decomposition behaviour of steels. Specimens were heat treated using a Theta Industries, Inc. Quench Dilatometer at 1200° C. for 3 minutes to dissolve the microalloying elements and achieve the desired large austenite grain size (˜100-120 μm), simulating the relatively course austenite grain size as produced by the twin roll strip casting process (i.e. CASTRIP® process). Samples were cooled at controlled, continuous cooling rates through the austenite (γ) to ferrite (═) transformation using helium gas as required. The transformation temperatures were taken at the changes in slope of the dilation curves. Optical metallography was used to determine the transformation products and CCT diagrams were constructed. Hardness measurements were taken using a Vickers hardness indenter with a 2.5 kg load.
TABLE-US-00001 TABLE I Chemical composition of UCS steels Element (Wt %) C-Mn V Nb Nb-V C 0.035 0.025 0.036 0.032 P 0.010 0.012 0.013 0.012 Mn 0.89 0.92 0.87 0.92 Si 0.23 0.22 0.21 0.22 S 0.002 0.001 0.002 0.003 Ni 0.067 0.046 0.039 0.042 Cr 0.061 0.051 0.043 0.037 Cu 0.12 0.12 0.08 0.13 Al 0.003 0.003 0.003 0.003 Nb 0.001 0.002 0.040 0.038 Ti 0.002 0.002 0.002 0.002 V 0.003 0.043 0.003 0.042 N 0.007 0.008 0.007 0.006
 Laboratory age hardening heat treatments were carried out using a continuous annealing simulator, using electric resistance for heating. Time-temperature thermal cycles were designed to simulate the thermal cycle utilized in annealing furnaces of a continuous hot dip galvanizing line, covering a range of peak temperatures and times. Results are compared to commercial hot dipped galvanized products the continuous hot dip galvanizing line used in commercial production.
 The CCT diagram for the base C--Mn, C--Mn--Nb and C--Mn--V steels are shown in FIGS. 3-6. For the C--Mn steel, FIG. 3, polygonal ferrite was the first phase to form at all cooling rates, through 100° C./s. Bainitic ferrite was first observed at a cooling rate of 10° C./s. As the cooling rate increased, polygonal ferrite digressed into quasi-polygonal ferrite, which was progressively replaced by bainite and the structure became finer and increasingly chaotic.
 Referring now to FIG. 4, results for the V steel were very similar to those for the C--Mn steel, although the transformation start temperatures were slightly higher and the hardness increased for the V steel.
 FIG. 5 shows the CCT diagram for the Nb steel and the effect of Nb on decreasing the transformation start temperature. Ferrite formation was completely suppressed at a cooling rate of 30° C./s with a microstructure that was 100% bainitic. Bainite also began forming at lower cooling rates compared to the C--Mn (FIG. 3) and V (FIG. 4) steels, beginning at a cooling rate of 3° C./s.
 FIG. 6 shows the CCT diagram for Nb--V steel. The diagram is similar to that for the Nb steel, with the hardness of the Nb--V steel at each cooling rate being higher.
 Typical microstructures for the four grades are shown in FIGS. 7(a)-(d), where the cooling rate was 30° C./s. FIGS. 7(a) and 7(b) show quasi-polygonal ferrite surrounded by the chaotic structure of bainitic ferrite in the base C--Mn and V UCS steels, while FIGS. 7(c) and 7(d) show completely bainitic structures in the Nb and Nb--V UCS steels. Due to continuous cooling, the bainite morphology ranges from coarser lath-like plates, transformed at higher temperatures, to fine needle-like structures, formed at lower temperatures.
 The Vickers hardness measurements for each steel grade are compared in FIG. 8. The hardness increased with increasing cooling rate as the softer polygonal ferrite was replaced by the harder and more dislocated bainitic phase. FIG. 8 shows that the addition of the microalloying elements increased the hardness compared to the base grade, in the order of V, Nb then Nb--V. This can be attributed to both solid solution hardening and microstructural hardening.
 The combined total of the increase in hardness of the Nb grade and the V grade, over that for the C--Mn grade, was similar, although higher, than the increase recorded for the Nb--V grade, indicating that the hardness increase from Nb and V were approximately additive. This is supported by the microstructural evidence depicted in FIGS. 7(a)-(d). The C--Mn and V grades showed similar microstructures so the majority of the hardening due to the V addition was attributed to solid solution hardening, but both the Nb and Nb--V grades showed similar microstructural hardening due to the effect of Nb on hardenability. Thus, the solid solution hardening from the V is added on top of the microstructural hardening and solid solution hardening from Nb in the Nb--V grade.
 Evaluation of the CCT diagrams demonstrated that: the addition of V provides little to no microstructural hardening, but does provide some solid solution hardening, that Nb increases the hardenability of the steel by both microstructural hardening and solid solution strengthening, and that the dual addition of V and Nb provided only a modest benefit to strengthening over Nb only microalloyed UCS. Further, TEM evaluation did not find any Nb precipitation, indicating that most of the Nb was in solid solution and available for hardenability. The solid solution strengthening from V is thought to be weak to moderate, so it is not clear if this mechanism can account for the full increases in hardnesses, generally 20 Hv5 across all cooling rates, observed in the CCT work. It is also possible V--N pairs (as opposed to V in singularity) could contribute to increased strengthening.
 FIG. 9 shows the average (Av) yield strength (YS) and tensile strength (TS) results for a range of Nb microalloyed UCS steels (0.015% to 0.080% Nb), along with results from a 0.04% V trial grade. As illustrated by FIG. 9, with increasing Nb additions, strength levels increased, but the V alloyed grade showed only a small increase in strength over that for the typical base composition. At a similar microalloying addition, 0.04%, the Nb UCS steel showed an increase in yield strength of about 50 MPa higher than the comparable V UCS grade.
 The bands shown for the yield and tensile strengths reflect that there are differences in Mn content (0.72% to 0.87% Mn) between the steels. The higher Mn steels achieved strength levels at the top of the band, while the lower Mn steels recorded strengths towards the bottom of the band. Overall, the results showed that a 415 MPa UCS product can be achieved using a small addition of Nb and a 450-480 MPa UCS product can be achieved with higher Nb levels.
 The yield strength results for the 0.04 wt % V grade are shown in more detail and compared to the plain low-carbon base steel as a function of the hot rolling reduction in FIG. 10. The yield strength results were within the upper range for the plain carbon steel, indicating that the vanadium addition only provided solid solution hardening and did not provide strengthening via microstructural effects similar to niobium in the hot rolled condition.
 The typical microstructures for the 0.04% V UCS steel and the plain C--Mn base UCS steel are shown in FIGS. 11(a) and 11(b), respectively. The plain C--Mn and 0.04% V UCS steels were produced with a high cooling rate (low coiling temperature) and 30% hot reduction to achieve a higher strength, 380 MPa UCS product. As shown in FIGS. 11(a)-(b), the microstructures of the two compositions are similar. Vanadium did not provide any noticeable increase in hardenability to suppress the transformation of allotriomorphic ferrite to allow bainite to transform, as per the 0.04% Nb steel (see, FIG. 12(c)). The cooling rate was sufficient, however, to suppress the progress of polygonal ferrite and acicular ferrite transformed instead within the austenite grain. Slow cooling will lead to a polygonal ferrite microstructure and a 280 MPa product, from the same chemistry.
 The strengthening effect of Nb was found to be due to microstructural hardening, as no precipitation was detected. Transmission electron microscopy examination of two Nb microalloyed UCS steels did not reveal evidence of Nb precipitation. Further examination of the Nb UCS steels using the atom probe tomography technique also did not reveal any evidence of Nb precipitation and indicated that Nb was retained in solid solution. As shown in FIGS. 12(a)-(d), the progressive strengthening of the microstructure was due to the progressive replacement of polygonal and grain boundary ferrite with acicular ferrite and bainite, due to the strong hardenability effect of Nb. As austenite grain boundaries are a lower energy site for nucleation than intra-granular nucleation of acicular ferrite at inclusions, bainite could be expected to progressively displace acicular ferrite with increasing Nb content. After the allotriomorphic ferrite was replaced with bainite, the progressive strength increment recorded in the Nb UCS steel with further increases in Nb content was most likely accounted for by the progressive refinement of the microstructure.
 The CCT study and the evaluation of the microstructure of commercially produced Nb microalloyed coils have both shown that Nb provides an increase in hardenability. Even relatively small additions of Nb were found to suppress allotriomorphic ferrite formation, allowing bainite to nucleate, providing substantial strengthening at cooling rates readily achievable during ROT cooling of thin strip (see, FIG. 12(a), 0.015% Nb). Previous investigations have found Nb levels of 0.008% provided sufficient increase in hardenability to provide strengthening through microstructural hardening. As previously discussed with respect to FIGS. 5 and 6, bainitic microstructures were observed at cooling rates of 30° C./s and above for the 0.04% Nb and 0.04% Nb/V steels. In addition, due to the strong solute drag effect of Nb, austenite recrystallisation is suppressed following in-line hot rolling, providing a relatively large, unrecrystallized austenite structure, further aiding the hardenability of Nb UCS steels. The effect of hot rolling on the austenite microstructure in Nb microalloyed UCS steels is shown in FIGS. 13(a)-(b). FIG. 13(a) shows unrecrystallized autenite grains in the Nb steel after 15% hot rolling reduction. FIG. 13(b) shows unrecrystallized autenite grains in the Nb steel after 35% hot rolling reduction.
 The regime of the processing conditions in the process for forming ultra-thin cast strip facilitates the retention of Nb in solid solution, which in turn provides microstructural strengthening. Moreover, most of the Nb remains available for subsequent precipitation by means of age hardening. The retention of Nb in solution can be attributed to the following factors:
 i) Rapid cooling rate during solidification that avoids the formation of large interdendritic Nb rich precipitates commonly formed in both thick and thin continuously cast slabs.
 ii) Single hot rolling pass of limited total reduction and the short time interval before ROT cooling starts, suppresses strain induced precipitation of Nb, which usually occurs during rolling of hot rolled strip.
 iii) The enhanced hardenability imparted by Nb in solid solution suppresses grain boundary ferrite formation, promoting the acicular ferrite and bainitic transformations. As such, transformation start temperatures are low enough to suppress Nb precipitation from occurring in a continuous cooling environment such as the ROT of the ultra-thin cast strip process.
 As previously discussed, the TEM and APT examinations of the Nb microalloyed UCS trial steels did not reveal any Nb precipitation in the hot rolled condition, and the APT results showed the Nb to be retained in solid solution. Thus, making it available for subsequent strengthening by an age hardening heat treatment.
 Short time heat treatments were carried out using an electric resistance heated continuous annealing simulator, utilizing a time-temperature cycle to simulate the annealing section of a continuous hot dip galvanizing line for a range of peak temperatures and hold times at the peak temperature. The response of the 0.04% V, 0.04% Nb and 0.04% Nb+0.04% V trial UCS steels to simulated age hardening heat treatments have been evaluated and the results are summarized in FIGS. 14 and 15. FIG. 14 shows the effect of the peak heat treatment temperature on the strength of three microalloyed UCS steels for a single dwell time at the peak temperature. As shown, the Nb and V steel has a tensile strength of at least 550 MPa after age hardening at a peak temperature of between 700 and 750° C. for between 9 and 11 seconds. More specifically, the dwell time at peak temperature may be 10 seconds. Still further, the tensile strength may be between 550 and 600 MPa. The difference in strength between the 0.04% V coiled at 570° C. and 610° C. is due to microstructural changes, the 610° C. coiling temperature produces much more of the softer polygonal ferrite. FIG. 15 presents the peak temperature that produced the highest strength as a function of the dwell time at the peak temperature.
 Based on the simulated age hardening data, we determined that the 0.04% V UCS grade showed a smaller age hardening increment compared to the 0.04% Nb grade but the peak aging temperature was higher. The dual Nb/V UCS steel fits with the above observation as only a small additional increment in peak strength was obtained compared with the 0.04% Nb steel, Importantly, however, the greater width of the temperature range for peak age hardening for the dual Nb/V microalloying system suggests such an alloy design would provide a broader processing window for applying an age hardening heat treatment in a continuous strip processing environment. The CCT diagrams confirm that adding V to the Nb grade only provided a small influence on the microstructure of the thermally processed steel.
 We have found that niobium provides strengthening in the as-hot rolled condition, predominately through microstructural hardening, via directly increasing hardenability and indirectly by suppressing austenite recrystallisation, retaining the coarse as-cast austenite grain size. Additionally, vanadium had little to no influence on microstructural hardening but contributes to increasing strength in as rolled steel via solid solution hardening. The retention of Nb and V in solid solution provides the capability to achieve strengthening from age hardening heat treatments as compared to that achieved in the hot band. The temperature range for peak age hardening were slightly different for Nb and V, providing sufficient overlap that a dual Nb+V alloy design provides a wider peak hardening temperature range.
 The results of extensive laboratory studies and plant scale trials has shown that microalloyed (Nb, V), low C and carbon equivalent, UCS steels produced by the CASTRIP® process offer the capability to achieve a range of thin, high strength (HSLA), sheet steel grades in the hot rolled and galvanized condition. This capability significantly extends the strength/thickness range for hot rolled strip products.
 While the invention has been illustrated and described in detail in the foregoing drawings and description, the same is to be considered as illustrative and not restrictive in character, it being understood that only illustrative embodiments thereof have been shown and described, and that all changes and modifications that come within the spirit of the invention described by the following claims are desired to be protected. Additional features of the invention will become apparent to those skilled in the art upon consideration of the description. Modifications may be made without departing from the spirit and scope of the invention.
Patent applications by Daniel Geoffrey Edelman, Brownsburg, IN US
Patent applications by Walter N. Blejde, Brownsburg, IN US
Patent applications by NUCOR CORPORATION
Patent applications in class Continuous casting
Patent applications in all subclasses Continuous casting