NANO-PRECIPITATION STRENGTHENED COLD-ROLLED BATCH ANNEALED HIGH STRENGTH LOW ALLOY STEEL SHEET

Abstract

A high strength low alloy steel sheet product produced by a continuous strip process. The sheet product comprises from 0.045 to 0.06 weight percent Carbon, from 0.75 to 1.2 weight percent Manganese, from 0.02 to 0.04 weight percent Aluminum, and from 0.075 to 0.12% weight percent Titanium and an Fe balance. The steel sheet product has been subjected to cold rolling and batch annealing to form a steel sheet product having a yield strength of at least 500 MPa, a substantially ferritic microstructure with nano TiC precipitates and a hole expansion ratio of at least 60%.

Description

CROSS-REFERENCE TO RELATED APPLICATION

This application claims the benefit of U.S. Provisional Patent Application No. 63/341,811 filed May 13, 2022, which is incorporated herein by reference.

BACKGROUND

The present disclosure generally relates to a high-strength low-alloy steel having a yield strength of at least 490 MPa. The disclosure includes a process for manufacturing a low carbon Ti-microalloyed of steel for the development of cold-rolled and batch-annealed high-strength high-formable steels. In some embodiments, the process includes a clean steel practice combined with hot rolling and batch annealing utilized to retain precipitate strengthening in the final fully processed steel and microstructural attributes that results in increased forming characteristics.

Development of high strength low alloy steels (HSLA) has redefined steel usage in various industries, for example, construction, machinery, automotive, agriculture and transportation. The HSLA steel provides engineering as well as economic benefits such as, weight reduction, improved and excellent weldability, reduced cost of engineering constructions, safety of components, increased payloads for transportation and enhanced fuel economy for passenger cars. The automotive industry has benefitted through advent of higher strength HSLA steels as this steel has provided for increased vehicular safety, enhanced fuel economy, and minimized CO2 emissions. Automotive manufacturers are now leaning towards use of greener HSLA steels in an effort to minimize CO2 emission. Companies are requesting that steel manufacturers provide carbon-neutral or carbon-minimal HSLA steels and the global steel industries are gearing towards innovating steel processing ideas and technologies that provide sustainable solutions to green steel production. In this evolving perspective, mini steel mills employing continuous strip production (CSP) technology have significant advantage in reducing CO2 emissions over discrete integrated steel mills that runs on blast furnace ironmaking technology. Moreover, CSP mills integrated with advanced steel processing technologies are competing with thick slab casting discrete steel mills not only in production throughput but also in offering advanced high strength grades of steels with lean alloying and least cost of production.

HSLA sheet steels with a yield strength of at least 490 MPa are increasingly demanded for many stamped parts as well as tubular forms of autobody applications. High yield strength HSLA steels have not been widely developed. In the past, the addition of vanadium (V) and/or niobium (Nb) microalloyed HSLA steels was investigated. A solid solution was strengthened as well as microalloy V, Nb strengthened sheet steels and mostly with production through thick slab casters and with heavy alloying. Furthermore, these developments also received less attention because of the emergence of dual-phase (DP) steels offering high tensile strengths and strain hardening capacity. However, DP steels are heavily alloyed and the ferrite yield strength is lower due to pre-yielding when compared with similar-tensile strength HSLA steels.

SUMMARY

The present disclosure includes a high strength low alloy steel sheet product produced by a continuous strip process. The sheet product comprises from 0.045 to 0.06 weight percent Carbon, from 0.75 to 1.2 weight percent Manganese, from 0.02 to 0.04 weight percent Aluminum, and from 0.075 to 0.12% weight percent Titanium and an Fe balance. Wherein the steel sheet product has been subjected to cold rolling and batch annealing to form a steel sheet product having a yield strength of at least 500 MPa, a substantially ferritic microstructure with nano TiC precipitates and a hole expansion ratio of at least 70%.

The present disclosure includes a method for producing a high strength low alloy steel sheet product comprising from 0.045 to 0.06 weight percent Carbon, from 0.75 to 1.2 weight percent Manganese, from 0.02 to 0.04 weight percent Aluminum, and from 0.075 to 0.12% weight percent Titanium and an Fe balance, a yield strength of at least 500 MPa, a substantially ferritic microstructure with nano TiC precipitates and a hole expansion ratio of at least 60%, comprising the steps of: continuously casting a steel slab approximately 55 mm-85 mm thick; maintaining a temperature of the steel slab; hot rolling the steel slab to a steel sheet at a finishing temperature of 875-950C; cooling the steel slab; coiling the steel sheet at a temperature of 600-675C; cold rolling the steel sheet to 40-75%; and hydrogen batch annealing the steel sheet at temperature of 600-650C to achieve a fully recrystallized ferritic microstructure with nano precipitates of TiC.

The present disclosure includes a high strength low alloy steel sheet product produced by a continuous strip process. The sheet product comprising from 0.045 to 0.06 weight percent Carbon, from 0.75 to 1.2 weight percent Manganese, from 0.02 to 0.04 weight percent Aluminum, from 0.075 to 0.12% weight percent Titanium, 0.04 weight percent Niobium and an Fe balance. Wherein the steel sheet product has been subjected to cold rolling and batch annealing to form a steel sheet product having a yield strength of at least 500 MPa, a substantially ferritic microstructure with nano TiC precipitates and a hole expansion ratio of at least 60%.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1 is a prior art graph showing the effect of Mn and carbon content on the r m values of cold rolled annealed steels.

FIG. 2 is a prior art graph showing the relationship between the state of austenite recrystallization during hot rolling and the transformed ferrite grain size.

FIG. 3 is a prior art graph showing the influence of microalloying on the recrystallization stop temperature in low carbon HSLA steels.

FIG. 4 is a graph showing the yield and tensile properties of cold worked samples of selected steels after simulated batch annealing.

FIG. 5 is a graph showing the variation in the strength properties of Ti- and Ti—Nb steels across the coil width after batch annealing.

FIGS. 6a and 6b are micrographs showing the optical microstructures of Ti steel and Ti—Nb steel after hot rolling.

FIGS. 7a through 7d show the optical and scanning electron micrographs of batch annealed, Ti-steel, and Ti—Nb steel showing fully recrystallized microstructures.

FIGS. 8a and 8b show inclusions found in the final fully annealed sheet steels revealing mainly globular oxide and oxy-sulfide inclusions.

FIGS. 9a through 9c are TEM images of hot rolled Ti—Nb steel showing coarse precipitates along grain boundaries and an overview of fine precipitates and some localized areas of precipitates in row arrangement.

FIGS. 10a and 10b show the chemical composition of precipitates in hot rolled Ti—Nb steel.

FIGS. 11a and 11b shows the fine nano-sized precipitates within ferrite grains and some along the grain boundaries.

FIGS. 12a through 12c are images showing the precipitates in hot rolled Ti-Steel as examined in thin foils, distribution of fine precipitates.

FIG. 13 shows HAADF image and EDS mapping of precipitates larger than 20 nm were Ti, C and S bearing.

FIG. 14 shows an HAADF image and EDS mapping of nano-sized precipitates in hot rolled Ti-steel indicating Ti as major element.

FIG. 15 shows a graph of EELS spectrum showing the presence of Ti and carbon only in a fine precipitate.

FIGS. 16a and 16b show images of fine precipitates in ferrite in Ti-steel and HAADF image and EDS mapping of precipitates in Ti-steel revealing presence of Ti only in fine precipitates less than 10 nm.

FIGS. 17a and 17b show TEM images of fully recrystallized microstructures of Ti and Ti—Nb steels.

FIGS. 18a and 18b show images of the microstructures of Ti-steel after simulated batch annealing at 630° C. and 700° C. revealing significant ferrite grain growth after recrystallization.

FIGS. 19a through 19c show a grain boundary mapping of Ti-steel and Ti—Nb steel constructed from orientation imaging from electron back scattered diffraction (EBSD).

FIGS. 20a and 20b show an ODF of batch annealed Ti-steel and Ti—Nb steel showing ODF of batch annealed Ti-steel and Ti—Nb steel showing increasing γ-fiber orientations.

FIGS. 21a through 21d are charts showing the precipitate size distribution in hot rolled and annealed Ti-steel.

DETAILED DESCRIPTION

Mechanical property requirements for HSLA 490 sheet steels are listed in Table 1 as required by various common sheet steel OEM specifications. Table 1 shows the tensile property requirements of HSLA 490 annealed sheet steels as required by various specifications.

	TABLE 1
		Tensile	%
	Yield	Strength	Elonga-
	Strength, MPa	MPa	tion
Specification	Grade	min	max	Min	min
SAE J2340	490 X	490	590	560	14
SAE J2340	490 Y	490	590	590	12
ASTM A 1008	HSLAS Grade 70	480		585	14
	Class 1
ASTM A 1008	HSLAS Grade 70	480		550	14
	Class 2
ASTM A 1008	HSLAS-F Grade 70	480		550	16

In some cases, a Ti-only alloy approach was utilized because very fine sized TiC can be precipitated in hot rolled coils giving a significant increase in ferrite yield strength. Ti is effective in scavenging nitrogen from steel and tying up with most of the carbon to form nano sized precipitates provided hot rolling parameters are properly controlled. The extent of precipitation strengthening can be significant and is approximated by a σ_ppt=B. (wt. % alloy). In some embodiments, B is an average 1500 MPa/wt. % Ti. Carbon was maintained below 0.06 wt. % as excessive carbon to that required for TiC formation will cause Fe-carbide precipitation during coiling and may deteriorate drawability and stretchability. Solid solution elements such as Mn, Si are significant ferrite strengtheners as outlined in Eq [1], as set forth below. In some embodiments, Mn additions were kept low as an excess of Mn introduces mobile dislocations in ferrite during rapid transformations after hot rolling through its influence on Ar₃ temperature and such mobile dislocations interfere with yield strength through pre-yielding. Mn also adversely influences r_m value in annealed sheet steel as shown in FIG. 1.

Total strengthening will thus be additive of ferrite grain size and precipitation strengthening as given below in Eq. [1] as deduced from a treatise of strengthening mechanisms in ferrite.

σ_y(MPα)=88+37mn'83Si+2900N_free+17d^−1/2+σ_ppt+σ_d  Eq. [1]

In some embodiments, d is grain diameter in mm, a σ_ppt is the precipitation strengthening, σ_d is dislocation strengthening.

FIG. 1 shows the effect of Mn and carbon content on the values of cold rolled annealed steels. Grain size and strength of annealed sheet steels have a direct inheritance from hot rolled steel as indicated in FIG. 2. FIG. 2 shows the relationship between state of austenite recrystallization during hot rolling and the transformed ferrite grain size. A finer fully recrystallized austenite grain results in fine ferrite grain after transformation. Since TiC has limited solubility in austenite finish rolling is carried out at relatively higher temperature than Nb-microalloyed steels. Austenite is refined through repeated recrystallization rolling above non-recrystallization temperature (TNR) which is lower than that of Nb-bearing steels as shown in FIG. 3 from similar steel. As seen from experimental data in FIG. 3, Ti influence on TNR diminishes in excess of 0.10 wt. % and the TNR can be estimated to be around 917° C. An accelerated cooling after finish rolling also induces many nucleation sites for precipitation as well as induces non-polygonal fine ferrite formation. Additionally, to examine the influence of small addition of Nb on the starting austenite grain size and therefore, the final ferrite grain size a second alloy with Nb addition was considered. FIG. 3 shows the effects of microalloying on the recrystallization stop temperature in low carbon HSLA steels (C≤0.07 wt. %).

The present disclosure describes a low carbon, lean Ti-bearing alloy to develop ferritic HSLA 490 sheet steel in cold rolled-batch annealed condition. Strengthening in final sheet steel was obtained primarily through ferrite grain size and TiC precipitation strengthening with moderate solid solution strengthening. Precipitation of TiC nano-precipitates in hot rolled steel was induced through control of the finishing deformation temperature, cooling rate and coiling temperature. A batch annealing cycle appropriate of the current steel was evolved through pre-simulations to determine a full recrystallization annealing temperature. Actual mill production of sheet steels of various thicknesses revealed HSLA 490 steels could be successfully developed with excellent elongation and strain hardening index using a lean Ti-only HSLA approach. Nb addition of up to 0.04wt. % to the chosen alloy resulted in similar properties but didn't enhance tensile or formability properties over those of Ti-steel. TEM studies of microstructure and precipitation revealed fully recrystallized very fine ferritic microstructure with homogenously distributed nano TiC precipitates of 1.5-8 nm in size with an average 3.16 nm. Precipitate size or distribution were not affected by the recrystallization anneal temperature. Hole expansion ratio values of more than 90% were obtained for 1.2 mm thick Ti-HSLA sheet steels. Nano-sized TiC precipitates, fine recrystallized grain size and excellent internal cleanliness contributed to the hole expansion ratio values. Formability as indicated through stretchability indicator, hole expansion ratio far exceeds that reported by similar strength sheet steels.

HSLA steels with yield strength of at least 490 MPa having excellent stretch forming capability will be described herein. In some embodiments, a predominantly ferritic Ti-microalloyed HSLA steel may include a higher yield strength as well as formability through nano-scale precipitation in ferrite. The nano scale precipitates when precipitated within ferrite not only increases strain hardening of ferrite but also resists recrystallization during annealing.

In some embodiments, a very low carbon C—Mn based ferritic steel with Ti-only microalloying was selected to develop a family of HSLA sheet steels with a minimum yield strength of 490 MPa meeting SAE J2340 CR 490X specifications. Batch annealing was utilized in place of a conventional continuous annealing approach. Batch annealing provides results in higher total elongation and higher drawability compared to continuous annealing, batch annealing can be performed at lower temperatures, i.e. just high enough to complete recrystallization of cold deformed structure without disturbing the precipitate state, low temperature annealing prevents precipitation coarsening and grain growth is avoided, batch annealing also retains excellent flatness of thin gauge sheet steels as compared to continuous annealing because of prevention of heat buckling. Batch annealing with optimized cycle may result in better cross width property uniformity compared with continuous annealing.

Strength development is based primarily on precipitation strengthening and control of ferrite grain size. Details of alloy design, hot rolling processing approach and annealing parameter optimization through studies of microstructure and precipitation evolution at each stage of processing are described herein.

Steelmaking and Processing

Chemistry of the HSLA 490 steel is shown in Table 2. A clone chemistry with small Nb addition (<0.04 wt. %) was also chosen to examine partially recrystallized austenite grains on the final ferrite grain size and properties. Carbon was restricted to less than 0.06 wt. % (i) to minimize grain boundary iron carbide formation and allow only sufficient amounts for micro-alloy carbide formation, (ii) to control Ti:C stoichiometry for nano scale precipitation of TiC precipitates. Table 2 shows the chemistry of steels selected in the current study (wt. % max).

TABLE 2
Steel	C	Mn	P	S	Si	Ti	N	Nb	CEIIW	Pcm
Ti	0.060	1.1	0.015	0.003	0.30	0.12	0.009	Res	0.25	0.13
Ti-Nb	0.060	1.1	0.015	0.003	0.30	0.12	0.009	0.04	0.25	0.13
CEIIW: C + Mn/6 + (Cr + Mo + V)/5 + (Cu + Ni)/15;
Pcm: C + Si/30 + (Mn + Cr + Cu)/20 + Ni/60 + Mo/15 + V/10 + 5B.

The steel was manufactured utilized an electric arc furnace ensuring least slag carryover and low residuals through suitable choice of scrap mix. Deoxidation practice was aimed at low total dissolved oxygen so that Ti is not lost as TiO₂. Calcium treatment of oxide inclusions was optimized to modify alumina as well as sulfide inclusions. Low nitrogen content Fe-alloys were used to result in low nitrogen in the melt. Lower nitrogen helps to control the amount of titanium lost to nitrogen and thus aids in enhancing the Ti:C stoichiometry. The heats were cast at continuous caster with suitable mold powder to result in good surface quality of hot rolled coils. The steels were cast in to slabs of 55-65 mm thickness and fed continuously to 6-stand hot strip processing mill through a 290 meter long tunnel furnace maintained at temperature to eject slabs at exit temperatures of 1120-1130° C. Thin slabs (≤65 mm), fast casting speeds (5-5.5m/min) and low soaking temperature (1125° C.) helped achieving finer starting austenite grains than in a thick slab casting unit.

The slabs were hot rolled to 2.2-3.3 mm thickness using a six-pass reduction schedule and finish hot deformation temperature of 900-925° C. The hot rolled strip was immediately accelerated cooled to a coiling temperature of 600-675° C. A cooling rate of more than 30° C./s was employed using super reinforced laminar cooling to result in non-polygonal ferrite grains with substructures if possible. Hot rolled samples were collected from 600 cm inside of the outer lap of coils after cooling to evaluate mechanical properties and precipitation.

Cold Rolling and Choice of Annealing Parameters

The hot rolled coils were cold rolled to 0.85-1.2 mm thickness (60-65% cold reduction) and then annealed in a batch annealing furnace. The annealing cycle to be used for these cold worked coils were initially determined using laboratory simulations of batch annealing. Strip samples prepared from cold rolled steels of both types were subjected to a programmed batch annealing cycle at various temperatures in a box furnace and tested after furnace cooling for mechanical properties and microstructure. FIG. 4 shows the variation in yield strengths as a function of simulated batch annealing temperature. From the simulation results it was evident that a batch annealing temperature window of 600-650° C. was appropriate in achieving yield strength of 490 MPa and above in both the steels. Batch annealing cycle was utilized to fully recover and completely recrystallize the cold worked ferrite grains.

FIG. 4 shows the yield and tensile properties of cold worked samples of selected steels after simulated batch annealing. Based on this simulation result cold rolled coils were then batch annealed in a hydrogen annealing furnace at temperatures between 600-650° C. After cooling, samples were collected from both steel grades from head and tail ends for mechanical and microstructural property evaluation.

Mechanical and Microstructural Properties Evaluation

Samples from batch annealed coils were collected from both head and tail ends for various property evaluations such as hardness, tensile properties and microstructures. Tensile samples across the width were also tested to check cross width property variation.

Microstructural Characterization

Full thickness section along rolling direction was examined in optical microscope (Leica DMI5000-M) as well as in scanning electron microscope (Hitachi SU3500) for microstructural features and cleanliness study. Metallographic samples were mechanically polished to 1 μm diamond paste suspension followed by surface treatment in a Hitachi IM4000 ion milling system for orientation imaging mapping. An area of 128 mm×100 mm with a step size of 0.3 μm was chosen for EBSD analysis for orientation imaging and texture evaluation in a Hitachi SU3500 SEM.

Formability Evaluations

Hole expansion ratio tests were done on annealed sheet samples from both steels using an Interlaken SP400 test equipment at AMT-Fadi LLC. For the hole expansion tests, five square coupons of 100 mm×100 mm size were cut from each of three locations across the width-quarter width, center and three-quarter width locations for examining property homogeneity. Holes of 10 mm diameter (d_o) was punched at center of each coupon for HER tests.

The test coupons were clamped between a holder and die with a clamping force of 100 kN. A clearance of 12±1% of nominal sheet thickness was adopted conforming to ISO 16630:2009(E) specification. A conical punch with 60° angle was pierced through the hole at a speed of 0.25 mm/s and the crack appearance during piercing was monitored using digital imaging system. The piercing was done at least after 30 mins after punching the hole. Diameter of the holes after crack appearance was measured and hole expansion ratio, λ was calculated as

$\lambda =\frac{\left(d_f-d_o\right)}{d_o}\times 100,$

where d_o and d_f are initial and final diameter of the hole respectively.

Precipitate Examination in TEM

Precipitation in hot rolled samples as well as final fully processed sheet samples were studied in Talos L120C and Talos 200X electron microscope equipped with EDS X-ray spectrometer and electron-energy-loss spectroscopy (EELS). Samples of dimension 10 mm×10 mm were used to prepare TEM foils and replicas using a precision cutter from the center areas of the sample parallel to rolling direction. For foils preparation, samples were thinned by careful mechanical grinding and polishing down to a thickness of around 80 mm. 3 mm diameter discs were punched from thinned sheet followed electro-polishing in electrolyte of 10% perchloric acid in methanol at −40C and 16V. Additionally, extraction replicas were prepared from the same samples for fine precipitation analysis.

Results

Tensile properties of batch annealed sheet steels of both steel types are shown in Table 3. Table 3 shows that both the sheet steels met minimum yield strength of 490 MPa in fully annealed condition. Excellent total elongation values were obtained. Ti-steel sheets showed slightly higher elongation values compared to Ti—Nb steel. Tensile properties were also outstandingly uniform from head to tail of the coils as well as across the width of the coils as indicated in FIG. 5. Table 3 shows the tensile properties of batch annealed coils.

TABLE 3
		Thickness,
		mm	Yield Strength,	Tensile Strength,
		SAE 490X	MPa	MPa	% El
Steel	Location	Spec	490-590	560 Min	14 Min	n (10-eu)
Ti	Head	0.85-1.2	506-536	571-619	20-22	0.13-0.16
	Tail	0.85-1.2	511-554	587-628	20-24	0.13-0.14
Ti + Nb	Head	0.85-1.2	512-546	596-616	19-23	0.14-0.15
	Tail	0.85-1.2	509-544	594-622	21-23	0.14-0.14

FIG. 5 shows the variation in strength properties of Ti- and Ti—Nb steels across the coil width after batch annealing.

Microstructure

FIGS. 6a and 6b show the through thickness microstructures of hot rolled Ti-steel and Ti—Nb steel along the rolling direction. Ferrite grains in Ti-steel indicated significant non-polygonalilty due to accelerated cooling from finishing temperature. Microstructure of Ti—Nb steel represented finer and elongated ferrite grains. Ferrite grains are slightly coarser (d_av˜4.2 mm) in Ti-steel compared to Ti—Nb steel (d_av˜3.2 mm). The elongated ferrite grain structure in Ti—Nb steel were indicative of partial non-recrystallization of austenite during finishing rolling whereas Ti-steel represented a fully recrystallized microstructure prior to onset of accelerated cooling.

FIGS. 6a and 6b show the optical microstructures of Ti steel and Ti—Nb steel after hot rolling. FIGS. 7a through 7d show the optical and scanning electron micrographs of batch annealed (a), (c) Ti-steel, and (b), (d) Ti—Nb steel showing fully recrystallized microstructures.

FIGS. 7a through 7d revealed microstructures of fully processed Ti-steel and Ti—Nb steel after batch annealing. Ferrite grains were much refined and nearly equiaxed in Ti-steel. Ti—Nb steel also represented a fully recrystallized microstructure and ferrite grains are slightly elongated. Scanning electron images of both steels indicated further clarity on the completion of recrystallization and onset of ferrite grain growth.

Cleanliness of the steels were evaluated through inclusion mapping in scanning electron microscope with EDS. Ternary diagram of oxide and sulfide inclusions were plotted and inclusion area fraction of various prominent inclusion types were measured. FIG. 8a shows typical inclusions present in both steels revealing only globular oxides and oxy-sulfide inclusions. SEM-EDS microanalysis of globular inclusions indicated these to be primarily CaO—Al₂O₃ and CaO—Al₂O₃—CaS inclusions. A typical SEM-EDS mapping of such inclusions is shown in FIG. 8 revealing elemental presence of Ca, Al, O and S. Summary of significant inclusion types are shown in Table 4 indicating excellent cleanliness.

Table 4 shows a summary of area fraction of various inclusions present in both steels as evaluated through SEM-EDS microanalysis.

	TABLE 4
	Area Fraction
Inclusion Type	Ti-Steel	Ti—Nb-Steel
Oxides (Ca-aluminate, Alumina,	1.56 × 10−5	3.45 × 10−5
Spinel etc.)
Sulfides (MnS, CaS, CaMnS	0	2.68 × 10−6
Oxy-sulfides (Ca-aluminate-CaS,	2.92 × 10−5	4.0 × 10−5
Ca-aluminate-CaMnS)

Precipitation Studies in Hot Rolled and Annealed Samples Ti—Nb Steel

In general, both coarse and fine nano-sized precipitates were seen in the hot rolled and annealed specimens. Most of the coarse particles in hot rolled sample were 10-80 nm in size with an average size of 31.2±1.7 nm and seen in low magnification in thin foil samples in FIG. 9(a). Coarse precipitates were seen mostly at grain boundaries as shown in FIG. 9a. FIG. 9b shows an overview of fine precipitates of less than 8 nm in size and were distributed homogenously within ferrite. Fine precipitates were also found in local areas as interphase precipitates with row arrangements (FIG. 9c). The shape of the fine precipitates was quasi-square and round.

FIG. 9a shows TEM images of hot rolled Ti—Nb steel showing coarse precipitates along grain boundaries. FIG. 9(b) is an overview of fine precipitates and (c) some localized areas of precipitates in row arrangement.

Chemical compositions of the ppts were analyzed using replica samples using EELS and EDS. EELS was performed by Digital Micrograph program from Gatan/Ametek and EDS was performed by Velox from Thermo-fisher Scientific. A majority of precipitates were less than 20 nm were Ti and Nb containing as shown in FIG. 10a. No nitrogen was detected in them. These were (Ti,Nb)C precipitates and possibly formed during hot rolling. Particles coarser than 20 nm were mainly Ti4C2S2 or TiN inclusions. High angle annular dark field image (HAADF) and EDS maps of particles less than 10 nm indicated Ti as major concentration (FIG. 10b).

FIGS. 10a and 10b show the chemical composition of precipitates in hot rolled Ti—Nb steel. Majority of ppts less than 20 nm (a) were Ti and Nb containing and no nitrogen was detected. (b) HAADF and EDS maps of precipitates less than 10 nm in size indicating Ti as major element.

Precipitate analysis in fully processed annealed sheet samples of Ti—Nb steel revealed similar size, distribution of precipitates. Coarse precipitates of 20-50 nm in sizes were Ti and Nb rich. Fine precipitates of 3-11 nm were found distributed homogenously in ferrite. Ti was the main element detected in the fine particles as seen in FIG. 11(b).

FIGS. 11a and 11b show (a) an overview of fine nano-sized precipitates within ferrite grains and some along the grain boundaries. (b) HAADF images and EDS mapping of precipitates 3-11 nm in size in fully annealed Ti—Nb steel.

Precipitates in Ti-Steel

FIGS. 12 (a)-(b) show the precipitate distribution in hot rolled thin foil sample. Precipitates were homogenously distributed within ferrite grains and some precipitation could be observed along grain boundaries. No row arrangement of precipitates or interphase precipitates were seen in Ti-steel. Coarse precipitates were fewer than Ti—Nb steel and size distribution was narrower between 20-30 nm.

These particles were mostly characterized as Ti, S bearing as shown in HAADF image and EDS mapping in FIG. 13 and possibly formed during casting. Fine nano-sized precipitates with size ranging from 1.5 nm to 8.8 nm were found distributed uniformly within ferrite grains with an average of 4.3 nm as shown in FIG. 11(c). In order to identify the fine precipitates extraction replica technique was used and identified by EELS. HAADF image and EDS mapping as shown in FIG. 14(a) and STEM image and EELS maps identify them all fine precipitates as TiC precipitates (FIG. 14(b)).

FIG. 12 shows an overview of precipitates in hot rolled Ti-Steel as examined in thin foils, (C) distribution of fine precipitates. FIG. 13 shows an HAADF image and EDS mapping of precipitates larger than 20 nm were Ti, C and S bearing.

FIG. 14 shows HAADF image and EDS mapping of nano-sized precipitates in hot rolled Ti-steel indicating Ti as major element. FIG. 15 shows EELS spectrum showing the presence of Ti and carbon only in a fine precipitate.

Precipitate analysis in fully processed annealed sheet samples of Ti-steel revealed very fine nano-sized precipitates of 3-11 nm in size distributed homogenously in ferrite. Ti was the main element detected in the fine particles as seen in FIG. 15(a). All precipitates below 15 nm were identified as TiC (FIG. 15b).

FIGS. 16a and 16b show (a) Fine precipitates in ferrite in Ti-steel (b) HAADF image and EDS mapping of precipitates in Ti-steel revealing presence of Ti only in fine precipitates less than 10 nm.

Formability Evaluation: Hole Expansion Ratio Results

Table 5 lists the hole expansion ratio (HER) values obtained for both the Ti-steel and Ti—Nb-steel in fully processed condition. HER data from various locations across the width were obtained and an average of five samples from each location are summarized. As indicated from the results, outstanding HER values were obtained for both the steels. The values are very uniform across the width of the coils. The HER values are indicative of excellent edge ductility or stretchability of the steel so developed and is expected to perform well during stamping operations. Table 5 shows the hole expansion ratio testing results performed on annealed HSLA sheet steels.

	TABLE 5
	Thick-	Yield	HER Values λ, %
	ness,	Strength,	Quarter	Center	Quarter
Steel	mm	MPa	width	width	Width
Ti	1.2	501	96.7 ± 5.0	108.0 ± 6.7	103.4 ± 7.4
Ti +	1.2	517	90.2 ± 12.7	91.9 ± 10.1	89.3 ± 7.2
Nb

Both Ti-only and Ti—Nb bearing HSLA steels of the selected lean alloy compositions successfully yielded a yield strength minimum of 490 MPa in cold rolled batch annealed condition. Nb microalloying did influence the austenite grain size after finishing deformation as could be seen from elongated ferrite grain structure (FIG. 6b) and presence of grain boundary (Ti,Nb)C precipitates (FIG. 9a). In contrast, Ti-steel represented a fully recrystallized microstructure after finish deformation rolling and was manifested in final ferrite grain structure. Non-recrystallization temperature, T_NR for Ti-steel was estimated much lower (900° C.) than the finish deformation temperature employed compared to Ti—Nb steel (967° C.) as estimated using experimental evaluations. There was no advantage of adding Nb in the final cold rolled and annealed microstructure as the grain size advantage in hot rolled steel was later marginalized after cold rolling and recrystallized annealing. Review of tensile properties in batch annealed sheet steels indicated Ti-steel indicated a marginal higher elongation values and better tensile properties across the width of the coils.

FIG. 17 shows the TEM images of fully recrystallized microstructures of Ti and (b) Ti—Nb steels. Hot rolling processing parameters were optimized for achieving maximum precipitation strengthening in the hot rolled coils. Batch annealing parameters after cold rolling were optimized to attain full recrystallization annealing before ferrite grain growth commences. Softening behavior of cold worked samples as studied through batch annealing simulation (FIG. 4) indicated an operable temperature window between 600-650° C. for the selected alloy composition. Beyond this temperature ferrite grain coarsening observed and a substantial loss in strength detected as shown in ferrite microstructures in FIG. 18. FIG. 18 shows the microstructures of Ti-steel after simulated batch annealing at (a) 630° C. and (b) 700° C. revealing significant ferrite grain growth after recrystallization.

Both steels represented similar softening behavior with temperature. Optical micrographs as well as transmission electron micrographs of annealed samples did indicate full recrystallization of ferrite grains. (FIG. 6, FIG. 17). No substructures were observed in either of the steels. Ti-steel manifested a slightly finer final grain structure than the Nb-added Ti-steel. Further studies of texture evaluation and grain boundary orientation of ferrite grains were carried out using electron back scattered diffraction (EBSD).

FIG. 19 shows the grain boundary mapping of (a) Ti-steel and (b) Ti—Nb steel constructed from orientation imaging indicating significant fraction of high angle grain boundary (HAGB) (˜79-83%) in both steels revealing near complete recrystallization in both steels. Blue lines represent high angle grain boundaries with critical misorientation of 15o or red lines representing low angle boundaries with misorientation of 2 to 15o. Grain size distributions as calculated using grain boundary mapping is shown in (c).

Grain size distribution as estimated from grain boundary mapping is plotted in FIG. 18(c) and an average ferrite grain size of 3.9 mm and 4.6 mm were estimated for Ti-steel and Ti—Nb steel respectively.

The annealing textures of both Ti-steel and Ti—Nb steel are represented in FIGS. 20a and 20b and showed similar recrystallization textures. Both steels revealed increasing γ-fiber with presence of some α-fiber as well. The γ-fiber showed significant intensity concentration or maximum at {111}<112> positions albeit non-uniform along the γ-fiber location. Ti-steel revealed relatively a stronger γ-fiber alignment compared to Ti—Nb steel. The ferrite grain boundary orientation indicated in FIG. 19 revealed a mixed grain structure which is probably the reason for not obtaining a sharp and uniform γ-fiber texture. Sharper γ-fiber texture is typical of extra deep drawing quality steels with a coarser and equiaxed ferrite grain structure. FIG. 20 shows the ODF of batch annealed Ti-steel and Ti—Nb steel showing increasing g-fiber orientations.

Precipitation studies revealed significant nano-sized TiC precipitates (1.5 nm to 8 nm) in both steels strongly contributing towards the yield strength increment. The precipitates that formed after hot rolling remained mostly intact in final batch annealed steels as the size and distribution of fine precipitate did not change during recrystallization annealing. In Ti-steel most precipitates were finer than 10 nm in size and lesser precipitates coarser than 10 nm were found. In contrast, Ti—Nb steel showed relatively coarser (Ti,Nb)C precipitates. FIG. 21 shows the precipitate evolution in hot rolled to cold rolled and annealed microstructure. Distribution and average size of precipitates were mostly similar. Average TiC precipitate size could be evaluated as 3.16 nm and 3.96 nm in Ti and Ti—Nb steels respectively. These nano precipitates could successfully contribute to the yield strength of the ferrite. Noting an effective Ti available for precipitation in Ti-steel and Ti—Nb steel as 0.067 wt. % and 0.058 wt. % (after nitrogen stabilization and Nb precipitation) the strengthening increment calculated using Eq. [1] can be 102 MPa and 87 MPa respectively.

Hole expansion ratio data obtained in the current batch annealed steel indicated superior stretchability or edge ductility suitable for most stamping operations. In general, both sheet steels indicated high HER values because of internal cleanliness, finer ferrite grain size and nano precipitates. Higher HER values obtained for Ti-alloyed HSLA steel was possibly due to absence of grain boundary coarser (Ti,Nb)C precipitates and relatively finer annealed ferrite grain structure than Nb added steel.

In some embodiments, hole expansion ratio values and tensile properties of the steel provides steel that can be utilized for autobody stamping applications requiring high stretch ductility.

Notwithstanding that the numerical ranges and parameters setting forth the broad scope of the invention are approximations, the numerical values set forth in the specific examples are reported as precisely as possible. Any numerical value, however, inherently contains certain errors necessarily resulting from the standard variation found in their respective testing measurements.

Any numerical range recited herein is intended to include all sub-ranges subsumed therein. For example, a range of “1 to 10” is intended to include all sub-ranges between (and including) the recited minimum value of 1 and the recited maximum value of 10, that is, having a minimum value equal to or greater than 1 and a maximum value of equal to or less than 10.

In this application, the use of the singular includes the plural and plural encompasses singular, unless specifically stated otherwise. In addition, in this application, the use of “or” means “and/or” unless specifically stated otherwise, even though “and/or” may be explicitly used in certain instances. In this application and the appended claims, the articles “a,” “an,” and “the” include plural referents unless expressly and unequivocally limited to one referent.

As used herein, “including,” “containing” and like terms are understood in the context of this application to be synonymous with “comprising” and are therefore open-ended and do not exclude the presence of additional undescribed or unrecited elements, materials, phases, or method steps. As used herein, “consisting of” is understood in the context of this application to exclude the presence of any unspecified element, material, phase, or method step. As used herein, “consisting essentially of” is understood in the context of this application to include the specified elements, materials, phases, or method steps, where applicable, and to also include any unspecified elements, materials, phases, or method steps that do not materially affect the basic or novel characteristics of the disclosure.

Claims

1. A high strength low alloy steel sheet product produced by a continuous strip process, the sheet product comprising from 0.045 to 0.06 weight percent Carbon, from 0.75 to 1.2 weight percent Manganese, from 0.02 to 0.04 weight percent Aluminum, and from 0.075 to 0.12% weight percent Titanium and an Fe balance, wherein the steel sheet product has been subjected to cold rolling and batch annealing to form a steel sheet product having a yield strength of at least 500 MPa, a substantially ferritic microstructure with nano TiC precipitates and a hole expansion ratio of at least 60%.

2. The high strength low alloy steel sheet product of claim 1, wherein the TiC precipitates are approximately 3.0 to 7.0 nm.

3. The high strength low alloy steel sheet product of claim 1, wherein the steel sheet comprises 0.06 weight percent Carbon, 1.1 weight percent Manganese and 0.12 weight percent Titanium.

4. The high strength low alloy steel sheet product of claim 1, wherein the steel sheet comprises a hole expansion ratio of at least 70%.

5. The high strength low alloy steel sheet product of claim 1 further comprising a tensile strength of at least 570 MPa.

6. The high strength low alloy steel sheet product of claim 1 further comprising less than 0.04 weight percent Niobium.

7. The high strength low alloy steel sheet product of claim 1 wherein the steel sheet is produced from a slab of approximately 55 mm to 85 mm thick.

8. The high strength low alloy steel sheet product of claim 1, wherein the steel sheet is produced with a coiling temperature of 600 to 675C.

9. The high strength low alloy steel sheet product of claim 1, wherein the steel sheet is produced by cold rolling to 40 to 75%.

10. The high strength low alloy steel sheet product of claim 1, wherein the steel sheet is produced hot rolling the slabs to a 2.2-3.3 mm thickness using a six-pass reduction schedule.

11. The high strength low alloy steel sheet product of claim 1 wherein the steel sheet is produced by hot rolling the steel slab at a hot rolling finishing temperature of 875 to 950C.

12. The high strength low alloy steel sheet product of claim 11, wherein the hot rolled strip is cooled at a cooling rate of at least 30° C./s to a coiling temperature of 600-675° C.

13. A method for producing a high strength low alloy steel sheet product comprising from 0.045 to 0.06 weight percent Carbon, from 0.75 to 1.2 weight percent Manganese, from 0.02 to 0.04 weight percent Aluminum, and from 0.075 to 0.12% weight percent Titanium and an Fe balance, a yield strength of at least 500 MPa, a substantially ferritic microstructure with nano TiC precipitates and a hole expansion ratio of at least 60%, comprising the steps of:

continuously casting a steel slab approximately 55 mm-85 mm thick;

maintaining a temperature of the steel slab;

hot rolling the steel slab to a steel sheet at a finishing temperature of 875-950C;

cooling the steel slab;

coiling the steel sheet at a temperature of 600-675C;

cold rolling the steel sheet to 40-75%; and

hydrogen batch annealing the steel sheet at temperature of 600-650C to achieve a fully recrystallized ferritic microstructure with nano precipitates of TiC.

14. The method for producing a high strength low alloy steel sheet product of claim 13, wherein the step of hot rolling includes hot rolling the slabs to a 2.2-3.3 mm thickness using a six-pass reduction schedule.

15. The method for producing a high strength low alloy steel sheet product of claim 13, wherein the step of cooing the steel slab includes a cooling rate of more than 30° C./s.

16. The method for producing a high strength low alloy steel sheet product of claim 13, wherein the steel sheet includes a tensile strength of at least 570 MPa.

17. The method for producing a high strength low alloy steel sheet product of claim 13, wherein the steel sheet includes a hole expansion ratio of least 70%.

18. A high strength low alloy steel sheet product produced by a continuous strip process, the sheet product comprising from 0.045 to 0.06 weight percent Carbon, from 0.75 to 1.2 weight percent Manganese, from 0.02 to 0.04 weight percent Aluminum, from 0.075 to weight percent Titanium, 0.04 weight percent Niobium and an Fe balance, wherein the steel sheet product has been subjected to cold rolling and batch annealing to form a steel sheet product having a yield strength of at least 500 MPa, a substantially ferritic microstructure with nano TiC precipitates and a hole expansion ratio of at least 60%.

19. The high strength low alloy steel sheet product of claim 18, wherein the TiC precipitates are approximately 3.0 to 7.0 nm.

20. The high strength low alloy steel sheet product of claim 18, wherein the steel sheet comprises 0.06 weight percent Carbon, 1.1 weight percent Manganese and 0.12 weight percent Titanium.

21. A high strength low alloy steel sheet product produced by a continuous strip process, the sheet product comprising from 0.045 to 0.06 weight percent Carbon, from 0.75 to 1.2 weight percent Manganese, from 0.02 to 0.04 weight percent Aluminum, and from 0.075 to 0.12% weight percent Titanium and an Fe balance, wherein the steel sheet product has been subjected to cold rolling and batch annealing to form a steel sheet product having a yield strength of at least 550 MPa, a substantially ferritic microstructure with nano TiC precipitates and a hole expansion ratio of at least 60%.

22. The high strength low alloy steel sheet product of claim 21, wherein the TiC precipitates are approximately 3.0 to 7.0 nm.

23. The high strength low alloy steel sheet product of claim 21, wherein the steel sheet comprises 0.06 weight percent Carbon, 1.1 weight percent Manganese and 0.12 weight percent Titanium.

24. The high strength low alloy steel sheet product of claim 21, wherein the steel sheet comprises a hole expansion ratio of at least 70%.