FIELD OF INVENTION:

The present invention relates to production of high strength Nano-precipitation strengthened steel plates with Ultimate tensile strength > 600 MPa. More particularly, the present invention is aimed towards design of alloy chemistry and thermo-mechanical process to produce and optimize size distribution of nano size precipitates in high strength formable grades of steel as well as in high strength low alloy micro alloyed steel with lower cost.

BACKGROUND ART:

The development of modern steels is essentially based on tailoring the microstructure to achieve the required properties. Keeping in view the potential applications in automobile and structural components, an extensive research work is currently in progress worldwide for increasing the level of microstructural refinement from micro to nano-level to achieve the desired combination of properties. Developing nano-sized precipitates is one of such routes through which good combination of properties can be achieved. It is already known that precipitates in steels play a very important role in achieving their desired mechanical properties. The chemistry, shape, size and distribution of these precipitates are important parameters which influence the material’s properties.The present work aimed towards design of alloy chemistry and thermo-mechanical process to produce and optimize size distribution of nano size precipitates in high strength formable grades of steel (Steel 1) as well as in high strength low alloy micro alloyed steel (Steel 2). With the successful generation of nano-sized precipitates (Average size <10 nm), it was possible to achieve improved properties (UTS ranging from607 MPa and 884 MPa and impact toughness ranging from 31 to 72 J at -40 0C in steels 1 and 2 respectively) in these steels than the existing steel grades being used for similar applications with lower cost.

In the prior art, a PCT application WO 2004/108970 discloses Nano-Precipitation Strengthened Ultra-High Strength Corrosion Resistant Structural Steels. The chemistry / composition includes incombination, by weight, about:0.1 to 0.5 % carbon, about 8 to 17% cobalt (Co), 0 to about10% nickel (Ni), about 6 to 12% chromium (Cr), less than about 1 % Si, less than about0.5% manganese (Mn), and less than about 0.15% copper (Cu), with additives selected fromthe group comprising about: less than 3% molybdenum (Mo). Less than 0.3% niobium (Nb),less than 0.8% Vanadium (V), less than 0.2% tantalum (Ta), less than 3 % tungsten(W), andcombinations thereof, with additional additives selected from the group comprising about:less than 0.2% titanium (Ti), less than 0.2% lanthanum (La) or other rare earth elements,less than 0.15 % Zirconium (Zr), less than 0.005% boron (B), and combinations thereof,impurities of less than about: 0.02% sulphur (S), 0,012% phosphorous (P), 0.015 % Oxygen(O) and 0.015%nitrogen (N), the remainder substantially iron (Fe), incidentalelements and other impurities. The alloy is strengthened by nanometer scale M2C carbideswithin a fine lath martensite matrix from which enhanced chemical partitioning of Cr to thesurface provides a stable oxide passivating film for corrosion resistance. The alloy, with UTSin excess of 280 Ksi, is useful for applications such as aircraft landing gear, machinery andtools used in hostile environments, and other applications wherein ultrahigh-strengthcorrosion resistant, structural steel alloys are desired.

In another prior art, an US Patent US 6746548 B2, discloses Triple-Phase Nano-Composite Steels. The carbon Steels of high performance are disclosed that contain a three-phase microstructureconsisting of grains of ferrite fused with grains that contain dislocated lath structures inwhich laths of martensite alternate with thin films of austenite. The microstructure can beformed by a unique method of austenization followed by multi-phase cooling in a mannerthat avoids bainite and pearlite formation and precipitation at phase interfaces. The desiredmicrostructure can be obtained by casting, heat treatment, on-line rolling, forging, and othercommon metallurgical processing procedures, and yields Superior combinations ofmechanical and corrosion properties.

OBJECT OF THE INVENTION:

The principle object of the present invention is to provide a Nano-precipitation strengthened steel composition (in wt.%) comprising of:C: 0.04 – 0.0.09, Si: 0.30 –0.60, Mn: 1.0–1.60, S: 0.009 – 0.015, P: 0.009-0.0.015, Ti: 0.10-0.20 or 0.015-0.02, Nb: trace or 0.05-0.07, V: trace or 0.05-0.07, Cr: trace or 0.15-0.20, Al: 0.02 – 0.03 and Balance-Fe

Another object of the present invention is to provide a hot deformation behavior at optimum working regions is carried out from 950 ºC - 1000 ºC is at strain rates of 0.01 s-1 to 0.03 s-1, while that of 860-910 ºC is at strain rates of 0.1 s-1 to 0.3 s-1.
Another object of the present invention is to provide hot deformation with hot rolling to 4 and 9 mm thickness plates and cooled in air.

Another object of the present invention is to provide nano-sized precipitates of average sizes 6-7 nm.

Another object of the present invention is to provide steel with yield strength 485-559 MPa, tensile strength 607-884 MPa, % elongation 24-28%.

Another object of the present invention is to steel with Charpy V-notch impact energy at -40 0C along transverse direction is greater than 30 Joules.

SUMMARY OF INVENTION:

Two laboratory heats were made in a 100 kg laboratory induction furnace and cast into 100 mm2ingots. The ingots were initially hot rolled to 20 mm plates in a two high experimental rolling mill at RDCIS, Ranchi after soaking at 12000C for 3 hours and a portion of the plates were used for initial characterization and hot deformation studies. The remaining plates were hot-rolled with the regulated parameters after soaking at 1175 0C for 45 minutes to the finish thickness of 4 and 9 mm for steels 1 and steel 2 respectively. The chemical compositions of the experimental steels are given in Table 1. Steel samples for chemical analysis was cut into the pieces of 35×35 mm2 cross section. The steel chemistry was analyzed by optical emission spectrometer (OES), Model-THERMO, ARL 3460.Isothermal hot compression tests were carried out in a computer controlled Dartec Servo-hydraulic UTM capable of strain rate variation from 10-2 to 10s-1 with a maximum load capacity of 100 kN and an attached furnace providing high temperature capability. For the purpose of the compression tests, cylindrical samples of height 15 mm with a diameter of 10 mm were subjected to strain rate of as low as 10-2 to high strain rate of10s-1. Each sample was deformed up to 0.69 true strains (˜50% engg. strain).
The tensile test specimens were made as per ASTM specification E8-04and the tests were performed as per the above standard, at room temperature in an INSTRON 1195 model mechanical test system. Longitudinal Transverse Charpy V-notch specimens were machined from the as-rolled plates so that the notch was perpendicular to the rolling direction. The geometry of the specimens was identical to that of the standard Charpy V-notch impact specimens except the thickness of the specimens used was machined to ~2.5 mm and 5 mm for steels 1 and 2 instead of 10 mm. Sub-sized specimens were used because of the thickness constraint of the plates available. The sub-sized impact specimens were tested at various temperatures from room temperature of 25°C to -40 °C) in a Tinius Olsen, USA make impact tester in accordance with ASTM E23-92 specification.

The microstructural studies of both the steels were carried out using Optical, Scanning and Transmission electron microscopes (TEM). The specimens for TEM were prepared by electro polishing, specimens thinned up to 70µm at a voltage of 20V, and the electrolyte being methanol + butoxy ethanol + perchloric acid mixed in the ratio of 6:3:1 for a period of 20-25 minutes.

Table 1: Chemical composition of the experimental steels Steel Grade C Mn Si S P Cr Ti Nb V Al
Steel 1 Ferritic 0.041 1.11 0.53 0.012 0.011 --- 0.18 --- --- 0.027
Steel 2 Ferrite-pearlite/
Bainite 0.087 1.51 0.38 0.009 0.007 0.2 0.018 0.06 0.05 0.01

BRIEF DESCRIPTION OF THE ACCOMPANYING DRAWINGS:
Fig.1 illustrates DTA curves in heating mode for (a) steel 1 (b) steel 2 in accordance with the present invention;
Fig. 2 illustrates Flow curves for different temperatures and strain rate (a) SR 0.01 s-1 (b) SR 0.1 s-1 (c) SR 1 s-1(d) SR 10 s-12 in accordance with the present invention;
Fig. 3 illustrates Deformed microstructures after compression at strain rate of 0.1 s-1 at different temperatures; (a) T = 850 ºC, (b) T = 900ºC, (c) T = 950 ºC, and (d) T = 1000 ºC2 in accordance with the present invention;
Fig. 4 illustrates (a) Strain rate sensitivity map at true strain 0.6, (b) processing map at true strain 0.6 unsafe processing areas enclosed within shaded red regions in accordance with the present invention;
Fig. 5 illustrates SE images of microstructural variation across different regions of processing mapsin accordance with the present invention;
Fig. 6 illustrates (a) Precipitate size distribution for deformation at high temperature (1050 ºC) and low strain rate (0.01 s-1) corresponding to optimum processing region in map (b) TEM bright field image showing nano-sized precipitatesin accordance with the present invention;
Fig. 7 illustrates Flow curves at different temperatures and strain rates (a) strain rate 0.01 s-1, (b) strain rate 0.1 s-1, (c) strain rate 1 s-1, (d) strain rate 10 s-1in accordance with the present invention;
Fig. 8 illustrates Microstructures of deformed specimens (a) T 850 ºC, strain rate 0.1 s-1 (b) T1000 ºC, strain rate 0.1 s-1 (c) T 850 ºC, strain rate 10 s-1 (d) T 1000 ºC, strain rate 10s-1in accordance with the present invention;
Fig. 9illustrates Strain rate sensitivity maps of log (strain rate) vs. temperature vs. strain rate sensitivity (m) at true strain (a) 0.35 (b) 0.65in accordance with the present invention;
Fig. 10 illustrates Processing map with unsafe processing regions enclosed within shaded red region for (a) true strain, 0.35 and (b) true strain, 0.65in accordance with the present invention;
Fig. 11 illustrates Processing map at true strain 0.65 with microstructural variations under different processing conditions. Precipitate distribution analysis for optimumprocessing regions (1050 ºC at low strain rates)in accordance with the present invention;
Fig. 12 illustrates (a) Precipitate size distribution for deformation at high temperature (1050 ºC) and low strain rate (0.01 s-1) corresponding to optimum processing region in map (b) TEM bright field image showing nano-sized precipitatesin accordance with the present invention;
Fig. 13 illustrates (a) Optical and (b) SEM micrographs showing primarily ferritic microstructure with equiaxed grainsin accordance with the present invention;
Fig. 14 illustrates TEM images (a) Bright field image showing precipitates (b) Corresponding DF image of (a) showing precipitates distributions of size <10 nmin accordance with the present invention;
Fig. 15 illustrates Precipitate size distribution histogram for steel 1in accordance with the present invention;
Fig. 16 illustrates (a) dual phase structure of steel 2 as seen in the optical micrograph, (b) SEimage showing equiaxed ferrite grains and irregular carbide distributionsin accordance with the present invention;
Fig. 17 illustrates TEM bright field images showing (a) and (b) Precipitate distribution in the matrixin accordance with the present invention;
Fig. 18 illustrates Precipitate size distribution histogram for steel 2in accordance with the present invention.

DETAILED DESCRIPTION:

Thermal transitions and differential thermal analysis
Differential thermal analysis was performed on both the steels to observe any major thermal transitions taking place in the samples and also to determine the critical transformation temperatures. The heating curves obtained for the samples are given below (Fig. 1)

Steel 1
The heating curve for Steel 1 shows two major transitions occurring near 760 ºC and just above 910 °C. ThermoCalc calculations using TCFE 7 database indicates the AC1 and AC3 temperatures to be 763.45 °C and 910.78 °C

Steel 2
The heating curve for Steel 2 displays a near continuous transition occurring between 750 °C and 840 °C. ThermoCalc calculations using TCFE 7 database indicate the AC1 and AC3 temperatures to be 690 °C and 852 °C respectively.

[2] Hot Deformation Behaviours
Isothermal hot compression tests were carried out in a computer controlled Dartec Servo-hydraulic UTM capable of strain rate variation from 10-3 to 10s-1 with a maximum load capacity of 100 kN and an attached furnace providing high temperature capability. Samples were coated with glass based lubricant that additionally served as a protective layer preventing oxidation at high temperature. Temperature measurements of the samples were achieved through K type thermocouple attached to the samples.

For the purpose of the compression tests, cylindrical samples of height 15 mm with a diameter of 10 mm were subjected to strain rate of as low as 10-2 to high strain rate of 10s-1. Each sample was deformed up to 0.69 true strain (˜50% engg. strain).

The samples were soaked at the deformation temperature (850 ºC – 1050 ºC) for five minutes before conducting the test. A maximum variation of 2 ºC from the specified deformation temperature was maintained by controlling the current supply to the furnace.
After deformation, the samples were quenched in water and sectioned parallel to the direction of compression along the diameter to study the microstructural changes in the samples during the deformation.

Nano-precipitate Strengthened Ferritic Steel (Steel 1)
Flow Curves
The stress strain behavior at different strain rates and different temperatures are given below in Fig. 2.
The following general observations can be made on the nature of the stress strain curves:
A recovery type behavior is shown by the material at 850 °C at all strain rates. Possibly, as was observed earlier in the heat treatments and from the austenitization temperature calculations (Ac3 = 910.78 °C), the presence of the ferritic structure with high stacking fault energy enables recovery type behavior.
At 900 °C the curve seemingly appears to show a work hardening type behavior marked by increasing stress values with strain.
The unusual hardness peak at 950 ºC seen in the heat treatments is shown by the material here also, represented by the highest strength at 950 ºC among all temperatures although a DRX type flow curve is typical of the deformation at 950 °C at all strain rates.
Deformation at higher temperatures (1000, 1050 ºC) results in DRX type behavior.
Shapes of the curves become similar to each other at high strain rates.
Microstructural analysis of the deformed samples correlates the flow behavior with the deformed/ recrystallized grains (Fig. 3).
Figure 3(a) corresponding to the recovery type behavior shown at 850 °C at 0.1 s-1 strain rate, indeed shows the deformed grains along the horizontal flow direction. Additionally, a bimodal distribution of the grains is seen, indicating some recrystallization occurring in the sample.
Figure 3(b), (c), (d) all showed equiaxed distribution of grains, with slightly increasing grain sizes in that order. Evidently dynamic recrystallization has occurred in all these samples, indicated by the equiaxed structure of the grains. Additionally, it is noteworthy that for the samples deformed at 1000 °C, a bimodal distribution of grains can be seen with some large grains, possibly due to grain growth after the initial recrystallization.
Strain Rate Sensitivity and Processing Maps
Strain rate sensitivity maps and processing maps are used for the investigation of dynamic mechanical properties such as formability, instability, efficiency of the deformation etc.
Here, variation of strain rate sensitivity (m) or the power dissipation efficiency (related to m) with change in strain rate and temperature are plotted in the form of a contour map. The system under consideration consists of a source of power (the hydraulic power pack), stores of power (platen, ram etc) and a dissipater of power (the work piece). The input power ((s.) ¯e ?) transmitted to brings about plastic deformation with the accompanying temperature increase (power component G) and microstructural changes (power component J). The power partitioning between the G and J component is determined by the parameter m which is the strain rate sensitivity of the material and is unity for an ideal linear dissipater.

The strain rate sensitivity (m) is represented as:
m = (?log(s))/(?log?(e ?)) .... (2)
The value of ‘m’ depends on the processing temperature (T), strain rate (e ?) and strain (e).In order to study the evolution of m with increase in strain, the strain rate sensitivity maps were plotted for strain of 0.60. Cubic spline interpolation was applied to increase the number of data points using the experimental values as knots. Using the m value from above equation and the strain rate and temperature, contour maps were plotted to get the following maps below [Fig. 4 (a)].
The efficiency of power dissipation through microstructural changes is determined by the dimensionless parameter ? given by,
? = J/Jmax= 2m/(m+1) .... (3)
Here Jmax represents the maximum efficiency of power dissipation from a linear dissipater (((s.) ¯e ?)/2).
Stable material flow determined by the stability criterion ? proposed by Ziegler, used in the Dynamic Materials Model, can be mapped as a function of strain rate and temperature to obtain an instability map. The stability criterion ? for stable flow is given by,
? = (? ln?(m/(m+1)))/(? ln?e ? ) + m > 0 .... (4)
The flow instability may manifest itself in form of adiabatic shear bands, flow localizations, dynamic strain ageing, mechanical twinning or flow rotations. Superimposition of power dissipation map and a flow stability map represents a processing map that can be used to identify stable flow regimes for safe processing of materials.
From Fig. 4 (a) and (b) it can be seen that the high strain rate sensitivity and high efficiency regions occur at low strain rates at temperatures beyond 1000 ºC. At high temperatures, easy movement of dislocations and less dislocation pileup (due to cross slip and climb) facilitate deformation. On a microstructural scale, good working efficiencies are obtained under conditions of dynamic recrystallization that occurs more readily at high temperatures and low strain rates. Unsafe working regions occur at high strain rates and low temperatures and hence working under these conditions should be avoided.
The green regions represent regions with good working efficiency, while the red regions represent regions with poor workability. However, it is also important to stay away from unsafe processing regions, represented by the shaded red regions overlapped over the efficiency map [Fig. 4 (b)].
Here, the steel shows two optimum working regions –
above 1010 °C at strain rates above 7 s-1.
1000 – 1020 °C at strain rates of less than 0.03 s-1
890-900 °C at strain rates of 0.1 s-1to 0.3 s-1.

Microstructural variation with deformation parameters
The microstructural variations across different regions of the processing map are represented in Fig. 5.
Deformation at low temperatures (850 ºC), where sample is not fully austenitized results in grains, elongated along the direction of deformation. A few large grains elongated along the deformation direction are also present in the microstructure. Evidently, the grains deformed along the deformation direction are yet to undergo recrystallization, which can be a consequence of the fact that the structure during deformation is ferritic as the steel is yet to undergo austenitization. The grain size also becomes finer at high strain rates (Fig. 5).
Deformation at higher temperatures in the austenitic region (950º C) results in fully recrystallized equiaxed grains. Though all grains are equiaxed, a range of grain sizes are present in the microstructure. Grain size does not seem to vary by much with strain rates. However, it is important to note that at the highest deformation temperatures (1050º C), the grains acquire an irregular shape instead of an equiaxed shape as seen in Fig. 5.

Precipitate distribution analysis for optimum processing regions (1050 ºC at low strain rates)
Using the processing map, the precipitate distribution of optimum processing regions, i.e. at temperatures of 1050º C and strain rates of 0.01 s-1 were determined using the extraction replica technique, whereby the quenched deformed samples were sliced along the center and carbon extraction replicas were prepared from them. Based on the data of twenty different regions, the precipitate distribution can be sorted as is represented in the histogram in Fig. 6.

Both fine and coarse precipitates (>100 nm) are present, although precipitates less than 50 nm constitute the largest fraction of these precipitates. Most of the strengthening is derived from these fine precipitates (<50 nm). Care should be taken to avoid processing at very high temperatures, where precipitate coarsening may occur.
Nano-Ferrito-bainitic Steel (Steel 2)
Flow Curves
The stress strain behavior at different temperatures and strain rates are represented in Fig. 7.
Following general observations can be drawn regarding the nature of the flow curves:
At high temperatures and low strain rates, the curves show a distinct peak, followed by steady state stress behavior typical of DRX.
At strain rates of 10 s-1, nature of curves at all temperature becomes similar in terms of the shape of the curves.
Deformation at 850 ºC shows a work hardening type behavior (steadily increasing stress). Similar observations have been made for other HSLA steels as well.
At high strain rates, the DRX peak becomes broad and shifts to larger deformation strains.
As expected, increasing strain rates and decreasing temperatures (high Z value) leads to high dislocation density and low dislocation mobility, increasing the deformation stress.
Microstructures of specimens sectioned along compression axis for samples deformed at 850º C [Fig. 8 (a), (c)] show a dual phase microstructure comprising of martensite with ferrite, with the shape of grains elongated along the flow direction (horizontal). Evidently, the microstructure during deformation at 850º C comprises of both ferrite and austenite, the latter of which transforms to martensite on quenching in water after deformation. This may help explain the shape of the flow curves for deformation at 850º C as seen in Fig. 7.
Specimens deformed at 1000º C show an equiaxed microstructure [(Fig. 8 (b), (d)], though the martensitic transformation makes the observation of the grain boundaries difficult. Nevertheless, the visible grain boundaries indicate equiaxed grains, thereby indicating occurrence of DRX during the deformation.

Strain Rate Sensitivity and Processing Maps
Strain rate sensitivity maps and processing maps are used for the investigation of dynamic mechanical properties such as formability, instability, efficiency of the deformation etc.
Here, variation of strain rate sensitivity (m) or the power dissipation efficiency (related to m) with change in strain rate and temperature are plotted in the form of a contour map. The system under consideration consists of a source of power (the hydraulic power pack), stores of power (platen, ram etc) and a dissipater of power (the work piece). The input power ((s.) ¯e ?) transmitted to brings about plastic deformation with the accompanying temperature increase (power component G) and microstructural changes (power component J). The power partitioning between the G and J component is determined by the parameter m which is the strain rate sensitivity of the material and is unity for an ideal linear dissipater.
The strain rate sensitivity (m) is represented as,
m = (?log(s))/(?log?(e ?)) .... (5)
The value of ‘m’ depends on the processing temperature (T), strain rate (e ?) and strain (e).
In order to study the evolution of m with increase in strain, the strain rate sensitivity maps were plotted for strains of 0.35 and 0.65. Cubic spline interpolation was applied to increase the number of data points using the experimental values as knots. Using the m value from above equation and the strain rate and temperature, contour maps were plotted to get the following set of maps below (Fig. 9).

High strain rate sensitivity regions occur at low strain rates and high temperatures (bottom right corner of map), while low strain rate sensitivity occurs at high strain rates and low temperature regions (top left corner). At high temperatures, easy dislocation movement and the lack of dislocation accumulation and pile up due to low strain rates (dislocation annihilation prevalent over dislocation generation) makes the material more responsive to variations in strain rate. The practical result of this is, as strain rate hardening forestalls unwanted flow localizations like necking, high strain rate sensitivity is beneficial in forming and shaping processes. Hence, regions with high strain rate sensitivity (T>10000C, e?<0.1s-1)are suitable for metal working operations.High strain rate sensitivity regions are green, while low strain rate sensitivity regions are represented in red colour in Fig. 9.
The efficiency of power dissipation through microstructural changes is determined by the dimensionless parameter ? given by,
? = J/Jmax= 2m/(m+1) .... (6)

Here, Jmax represents the maximum efficiency of power dissipation from a linear dissipater (((s.) ¯e ?)/2).
Stable material flow determined by the stability criterion ? proposed by Ziegler,used in the Dynamic Materials Model, can be mapped as a function of strain rate and temperature to obtain an instability map. The stability criterion ? for stable flow is given by,
? = (? ln?(m/(m+1)))/(? ln?e ? ) + m > 0 .... (7)
The flow instability may manifest itself in form of adiabatic shear bands, flow localizations, dynamic strain ageing, mechanical twinning or flow rotations. Superimposition of power dissipation map and a flow stability map represents a processing map that can be used to identify stable flow regimes for safe processing of materials.
The processing maps calculated for strains of 0.35 and 0.65 are represented in Fig. 10 below.

High temperature and low strain rates are regions of high efficiencies. The efficiency exceeds 0.3 in such regions (bottom right corner). High efficiencies, in low stacking fault energy materials (austenite) occur in the dynamic recrystallization domain. A look at the flow curves for strain rates of 0.01s-1 and 0.1s-1 at 1050 and 1000 ºC verifies this speculation, as the stated strains are near the DRX peak in the flow curve of the above cases. However, a lack of efficiency exceeding 60% indicates super plastic deformation does not occur in the material.
Unsafe processing regions are present within the shaded red regions in the maps. For the large strain deformations, the instability region at high strain rates and low temperatures can be the result of shear banding etc. The instability region at low strain rates and between 900-1000º C may appear on account of change in precipitate characteristics during deformation.
Overall, the optimum processing regions for working of this steel are in the range of:
Above 1000º C at strain rates of 0.01 s-1 to 0.03 s-1.
Above 1020º C at strain rates exceeding 1 s-1.
Microstructural variation with deformation parameters
The microstructural variations under different processing conditions for a true strain of 0.65 are represented in the Fig. 11.
It is obvious that the steel on deformation at 850º C is not austenitized completely. Thus the deformation of a dual phase structure of austenite and ferrite takes place. While some recrystallization may take place in the austenite, at high strain rates, the grains acquire an elongated structure in direction of deformation pointing to a lack of recrystalliztion at these strain rates (See Fig. 11).
At higher temperatures (950º C, 1050º C) of deformation irrespective of strain rate, the structure acquires an equiaxed configuration and lacks any obvious deformation structure, indicating occurrence of recrystallization (Fig. 11).

From the processing map, choosing the region of high strain rate sensitivity i.e. at high temperatures and low strain rates, extraction replicas were prepared for the analysis of their precipitate distribution. Based on the data of twenty different regions from the extraction replicas, the distribution of precipitate is plotted in the histogram in Fig. 12.
While the histogram shows maximum number of precipitates belonging to the 20-40 nm range, a substantial number of precipitates are below 20 nm size as well. It is important to note here that processing at excessively high temperatures for a long time e may coarsen some of the precipitates making them lose their effectiveness.
[3] Finish Rolling of the Steels
Steels 1 and 2 are finish-rolled to 4 and 9 mm thickness respectively in the experimental rolling mill at RDCIS.
[4] Microstructures of the Finish-Rolled Steels
The as-rolled steels were characterized for their microstructure by optical, SEM and TEM technique. While the base microstructure and predominant phases were easily characterized by optical and SEM, precipitate studies required TEM bright field images. For the purpose of TEM sample preparation, discs of 100 micron thickness with a diameter of 3 mm were subjected to precision ion polishing (PIPS) after dimpling the discs in the centre to reduce their thickness. All the TEM characterizations were done in a Technai F 30 TEM.
Nano- precipitated ferritic steel
The SEM images in SE mode of the as-received sample (Fig. 13) shows equiaxed ferritic grains, indicating a recrystallized structure with no obvious signs of a retained history of deformation. At higher magnifications however, some coarse precipitates can be observed in the structure.
TEM images (Fig. 14) show a uniform distribution of precipitates in the matrix with a substantial precipitate density.
The histogram in Fig. 15 represents the distribution of precipitate sizes in the observed sample.
Median size of the precipitate is 6.8 nm, while the average is 6.3 nm with a standard deviation of 2 nm. Like the previous case, the smallest precipitates in this case measured less than 2 nm as well.

Nano- precipitated ferrite-bainitic steel
Optical analysis of the second composition shows a dual phase structure with a dark phase (presumably some carbides) and a light phase (ferrite) (Fig. 16(a)). Though the grains are equiaxed (meaning recrystallized), there is substantial variation in the grain size in the sample.
The SE images show that while ferritic grains are equiaxed, the carbide distributions are irregular with some bainitic type appearance (Fig 16 (b))
Further analyses of the precipitates were done using bright field mode of the TEM. Most of the observed precipitates were in the range of 5-20 nm in size. Some dislocation precipitate interaction can be seen in the images as well (Fig. 17).
The size variation of the precipitates is represented in the histogram in Fig. 18.
Median size of the precipitates was found to be 7 nm which is about the same as the average value with a standard deviation of 2 nm.

[5] Mechanical Properties
The tensile and impact toughness properties (using sub sized specimens) of as received steel samples are presented in Table 2 below:

Table 2: Mechanical Properties

Grade of Steel Charpy V-notch Energy (J)* YS (MPa) UTS (MPa) % El
RT 0° C -20° C -40° C
Steel 1 100 72 60 31 485 607 24
Steel 2 92 84 84 72 559 884 28

Numerous characteristics and advantages of the invention covered by this document will be set forth in the foregoing description. It will be understood, however, that this disclosure is, in many respects, only illustrative. Changes may be made in details, particularly in matters of shape, size, and arrangement of parts without exceeding the scope of the invention

We Claim

1. A Nano-precipitation strengthened steel composition (in wt.%) with hot deformation and air cooling comprising essentially of:
C: 0.04 – 0.0.09, Si: 0.30 –0.60, Mn: 1.0–1.60, S: 0.009 – 0.015, P: 0.009-0.0.015, Ti: 0.10-0.20 or 0.015-0.02, Nb: trace or 0.05-0.07, V: trace or 0.05-0.07, Cr: trace or 0.15-0.20,
Al: 0.02 – 0.03 and Balance-Fe

2. The Nano-precipitation strengthened steel as claimed in claim 1, wherein the hot deformation behavior at optimum working regions is carried out from 950? C - 1000? C is at strain rates of 0.01 s-1 to 0.03 s-1, while that of 860-910? C is at strain rates of 0.1 s-1to 0.3 s-1.

3. The Nano-precipitation strengthened steel as claimed in claim 2, wherein the hot deformation is done with hot rolling to 4 and 9 mm thickness plates and cooled in air.

4. The Nano-precipitation strengthened steel as claimed in claim 3, wherein thesaid thickness can be either ferritic microstructure in or ferrito-pearlitic/bainitic microstructure.

5. The Nano-precipitation strengthened steel as claimed in claim 3, wherein the said thickness has nano-sized precipitates of average sizes 6-7 nm.

6. The Nano-precipitation strengthened steel as claimed in claim 2, wherein the said thickness of steel possess an yield strength 485-559 MPa, tensile strength 607-884 MPa, % elongation 24-28%.

7. The Nano-precipitation strengthened steel as claimed in claim 2, wherein the said thickness of steel possesses an Charpy V-notch impact energy at -400 C along transverse direction is greater than 30 Joules.