Effects of Temperatures of Rolling and Annealing on Microstructures and Tensile Properties of Low Carbon Ferritic Low Density Steels

Abstract

The microstructural modification and the effect of cold/warm rolling and annealing on tensile properties of ferritic low density steels containing different amounts of carbon (0.0035 mass%, Steel 1 and 0.04 mass%, Steel 2) and aluminium (6.8 mass%, Steel 1 and 9.7 mass%, Steel 2) were investigated with respect to amount of ferrite, grain size and formation of Fe₃Al intermetallic. The Steel 1 was deformed by cold rolling at room temperature and Steel 2 was deformed through warm rolling at 250°C to avoid prior cracks formation. After rolling, samples of Steels 1 and 2 were annealed at 900°C for 5 minutes. The yield strength, ultimate tensile strength and ductility increase while yield ratio decreases with annealing. This is due to formation of Fe₃Al intermetallic and reduction of ferrite grain size. Due to high interstitial Carbon in Steel 2, yield point phenomenon was observed which may reduce the formability of the steel. The cracks started during tensile test at the ferrite grains and ferrite grain boundaries and the fracture took place in the quasi cleavage mode in Steels 1 and 2. The grain refinement and the microstructural features due to rolling and annealing are accountable for the good combination of strength and ductility of the both steels.

1. Introduction

Addition of 1 mass% Al reduces the density of steels by 1.3% compared to conventional steels.^1,2,3) This has led to the development of Fe–Al–C low density automotive steels having high specific strength in order to improve the fuel efficiency and hence minimise the hazardous carbon dioxide gas emission. Addition of Al also stabilises the ferrite (BCC) phase and stimulates the precipitation of Fe₃Al intermetallic during cooling, below 552°C in Fe–Al low density steels.^4,5,6,7) On the basis of the phases, low density steels can be categorized into three types: single phase (only ferrite or austenite), two phases (ferrite or austenite and Fe₃Al intermetallic ) and multiphase (ferrite, austenite and Fe₃Al intermetallic).^8,9,10,11) Tensile properties are strongly dependent on these phases, for example the ferrite based low density steels with C(<0.05)-Al(5-9)-Mn(<5) mass% have ultimate tensile strength (UTS) and total elongation ranging from 200 to 600 MPa and 10 to 40% respectively.⁷⁾ In ferritic low density steels, a higher content of Al leads to an increase in the strength and a decrease in the elongation and formability due to the formation of Fe₃Al intermetallic with relatively higher hardness.

In the past, a number of studies on cold rolling and annealing of low-carbon low density steels have been reported.^12,13,14,15) As for example, Rana et al.⁹⁾ studied the crystallographic texture in cold rolled and annealed (at 900°C for 1 minute) steels having Al contents between 6.8 to 9.7 mass%. However, cold rolling and warm rolling of low density steels has not been widely reported and discussed. The reports on the fracture behaviour of the low density steels are sporadic. Sohn et al.¹⁶⁾ studied the cleavage cracking in low density steels containing 0.3C-3Mn-(4 to 6)Al mass% and showed that crack initiation takes place at κ-carbides at the ferrite grain boundaries. The crack propagated through the ferrite grains following cleavage fracture. Sohn et al.¹⁷⁾ highlighted the importance of annealing in improving the tensile strength and ductility by changing the fracture mechanism to mixed ductile and quasi cleavage fracture from cleavage fracture. However, very limited study has been reported on improving the strength and the effect of Fe₃Al intermetallic, through rolling and annealing processes. Therefore, investigations attempting to improve the ductility of high Al low density steels by modification of the microstructure through rolling and annealing are desired.

The present study aims to underpin the correlation between the compositions, microstructure development during cold and warm rolling and subsequent annealing with the deformation behaviour to design the processing stages for optimisation of microstructures and the tensile properties. In this study two steels with different carbon (0.0035 mass%, Steel 1 and 0.04 mass%, Steel 2) and aluminium (6.8 mass%, Steel 1 and 9.7 mass%, Steel 2) contents were examined. The cold and warm rolling and annealing treatments are designed to achieve ferrite grain refinement. The effects of chemical composition of the steels and the rolling and annealing parameters on the tensile properties were investigated. The tensile deformation and fracture of the steels were linked with the microstructures.

2. Experimental Procedure

The nominal chemical compositions of the low density steels, investigated in the present study and the density of steels, estimated by Archimedes’ principle are given in Table 1. The Thermo-Calc software with database TCFE7, was used to predict the equilibrium phases at room temperature. The specimens of Steel 1 of 25 mm thickness were reheated at 1250°C for 1 h and hot rolled at 900°C (finish rolling temperature (FRT)) to 3 mm thickness, air cooled to room temperature and then cold rolled to 1.2 mm thickness using multiple passes. The cold rolled specimens of Steel 1 were annealed at 900°C for 5 minutes. The hot rolling was performed on the specimens of Steel 2 at 900°C to 3.9 mm thickness. The hot rolled specimens were cooled to room temperature in air. After that, the specimens were heated to 300°C for 30 minutes and then subjected to warm rolling (rolling temperature ~250°C) to the thickness of 1.2 mm. The warm rolled specimens were annealed at 900°C for 5 minutes. The rolling and annealing schedules of Steel 1 and Steel 2 are summarised in Fig. 1. Moreover, cold rolling was given to a few strips of both the steels after hot rolling to examine if they can be directly cold rolled and thus to obtain insight about their cold rollability in order to determine the strategy of processing them to sheet form.

Table 1. Chemical compositions and density estimated by Archimedes’ principle of Steel 1 and Steel 2.

Steel name	C (mass%)	Al (mass%)	Ti (mass%)	Fe (mass%)	Density (g/cm3)	Density reduction relative to Fe
Steel 1	0.0035	6.8	0.10	balance	7.14	8.46%
Steel 2	0.04	9.7	0.09	balance	6.81	12.70%

Fig. 1.

Schematic diagram of rolling path and heat treatment conditions of (a) Steel 1, (b) Steel 2, and (c) Sample indicating rolling direction (RD), transverse direction (TD), and normal direction (ND). RT: Rolling temperature, FRT: Finish rolling temperature. (Online version in color.)

The microstructure analyses were done through ImageJ software on the surface from the rolling and normal directions of the strips (Fig. 1(c)). The specimens were polished to 0.05 μm SiO₂ finish and then etched with 5% nital solution (5 ml concentrated HNO₃ in 95 ml ethanol) and examined on a Leica DM 2500M optical microscope and in quanta 450 field emission gun scanning electron microscope (FEG-SEM). The compositions of the precipitates and the matrix were determined using energy dispersive X-ray spectroscopy (EDS). The operating voltage of 20 kV and working distance of 10 mm were utilised in the FEG-SEM. X-ray diffraction was done in the diffraction angle (2θ) range of 40–120° (Co-K_α target, scan step 0.02° and scan speed 0.5 per minute) to confirm the phases present at room temperature in the hot rolled and annealed samples.

Tension tests were carried out on the specimens of both the studied steels at a strain rate of 1×10⁻³ s⁻¹ at room temperature using a servo-electric universal testing machine of ±100 kN capacity (INSTRON 8862). Tests until complete fracture were carried out under strain control mode with an extensometer (capacity: +50% to −10%). A ASTM E-8 specimen geometry with 25 mm gauge length and 6 mm width was used. The tests were controlled and data acquisitions were performed using Instron Bluehill tensile test software available with the test system. Then fracture surfaces were observed in the FEG-SEM. Vickers microhardness (load: 100 g) and macrohardness (load: 10 kg) measurements on the steels at five random spots were done and the average values are reported.

3. Results

3.1. Thermodynamic Calculation of Solidification Sequences and Equilibrium Phases at Room Temperature

Figures 2(a) and 2(b) show the variations of phase contents with temperatures in Steel 1 and Steel 2 calculated using Thermo-Calc software. The solidification in Steel 1 (Fig. 2(a)) starts with the formation of ferrite at 1537°C and completes as ferrite at 1532°C. The solidification range of the alloy is 5°C. The Ti carbide phase with the face centred cubic (FCC) crystal structure is expected to form at 1150°C. With further cooling, Laves phase, containing Fe and Ti, with hexagonal crystal structure forms at 250°C and is stable up to room temperature. At 100°C, the equilibrium volume fractions of the phases are: ferrite ~ 0.9988, Ti rich carbide ~ 0.0003 and Lave phases ~ 0.00081. Solidification in Steel 2 (Fig. 2(b)) begins at 1522°C as ferrite and ends at 1492°C as ferrite and solidification range of the steel is 30°C. In Steel 2, the Ti-rich carbide phase (FCC crystal structure) nucleates at 1350°C. The higher carbon in Steel 2 raises the temperature of formation of Ti-rich carbide phase. Kappa phase with iron, aluminium, carbon and titanium forms at 900°C. At 100°C, the stable phases and their volume fraction are: ferrite ~ 0.99491, Ti rich carbide ~ 0.00158 and kappa phase ~ 0.00351. The solidification sequence and the temperature ranges of both the steels are given in Table 2. Thermo-Calc predicted compositions of the first solid ferrite and the last liquid of Steels 1 and 2 presented in the Table 3 show different concentrations of the elements in the ferrite and in the liquid or micro-segregation. The trend of micro-segregation is diverse for different elements and can be quantified by using the partition ratio. The partition ratios of the elements were calculated after taking the ratio of the composition of the last liquid to the first solid (ferrite) for the major solute elements. The partition ratio is greater than 1 thereby indicating that the first ferrite to form during solidification gets depleted with these solutes which are rejected to liquid. Table 3 shows that the first solid ferrite of Steels 1 and 2 is depleted primarily of solute elements C and Ti. The Al, having the lowest partition ratio, is nearly uniformly distributed in ferrite and the last liquid in Steel 1 and Steel 2.

Fig. 2.

The thermodynamics calculation of (a) Steel 1 and (b) Steel 2. (c) Variation of density of the Steel 1 and Steel 2, estimated using Thermo-Calc with temperatures. (Online version in color.)

Table 2. Solidification sequence of the steels.

Steel name	Solidification start temperature (°C)	Solidification finish temperature (°C)	Solidification range (°C)	Solidification sequence
Steel 1	1537	1532	5	Liquid→Liquid+Ferrite→Ferrite
Steel 2	1522	1492	30	Liquid→Liquid+Ferrite→Ferrite

Table 3. Thermo-Calc predicted composition (mass fraction) of first solid and last liquid in Steel 1 and Steel 2.

	Elements	First solid (ferrite)	Last liquid	Partition ratio
Steel 1	C	0.000006	0.00024	40
	Al	0.06618	0.07158	1.08
	Ti	0.00055	0.00247	4.49
Steel 2	C	0.00029	0.00315	10.86
	Al	0.09717	0.10288	1.058
	Ti	0.00084	0.00240	2.85

Figure 2(c) shows the variation of density of both steels with temperature predicted by Thermo-Calc. In Steel 1 and Steel 2 the density increases rapidly with the progress of the solidification when the stable phase ferrite with body centred cubic (BCC) crystal structure is present and it is associated with rapid volume contraction during solidification changing the solute distribution between solid and liquid, listed in Table 3. The rate of increase of density decreases sharply at 1150°C in Steel 1 and at 1350°C in Steel 2 due to the precipitation of Ti rich carbides with FCC crystal structure having higher specific volume resulting in lower effective volume contraction. As a consequence, with further cooling, the density of the Steels 1 and 2 continues to rise slowly due to gradual volume contraction. The measured room temperature densities of Steel 1 and Steel 2 are 7.14 gm/cm³ and 6.81 gm/cm³ (Table 1) respectively.

3.2. Microstructures in Different Stages of Processing of Hot, Cold or Warm Rolling and Annealing

Microstructures of hot rolled samples of Steel 1 (Fig. 3(a)) and Steel 2 (Fig. 3(b)) consist of large grains of ferrite. Microstructures of Steel 1 after 60% cold rolling and annealing at 900°C for 5 minutes (Fig. 3(c)) and after 69% warm rolling (rolling temperature 250°C) and annealing at 900°C for 5 minutes of Steel 2 (Fig. 3(d)) exhibit a substantial reduction in ferrite grains size, given in Table 4. The microstructures of cold rolled and annealed sample of Steel 1 (Fig. 3(e)) and warm rolled and annealed sample of Steel 2 (Fig. 3(f)) show the presence of Fe₃Al intermetallic in the ferrite grains and along grain boundaries. The area fraction of Fe₃Al intermetallic is relatively larger in Steel 2 than in Steel 1. The area fractions of ferrite and Fe₃Al intermetallic of Steel 1 and Steel 2 for different processing conditions are measured after taking average of area fraction using 10 continuous fields of views of microstructures, given in Table 5. The EDS line profiles in Steel 1 (Fig. 4(a)) and Steel 2 (Fig. 4(b)) after annealing show that the Fe and Al profiles are uniform in Steel 1, whereas in Steel 2, small fluctuations of Fe and Al are observed along the grain boundary. Moreover, Ti-rich precipitate peaks can be observed in EDS profile along grain boundaries of Steel 1 (Fig. 4(c)). The counts of Al atom in Steel 1 is relatively lower than in Steel 2 and the fraction of Fe₃Al intermetallic can be larger in annealed sample of Steel 2. The XRD plots, shown in Fig. 5, confirm that at the room temperature, only ferrite (BCC) peaks appear in both steels before and after rolling and annealing.

Fig. 3.

The 5% nital etched microstructures of Steel 1 in (a) hot rolled (optical), (c) and (e) cold rolled and annealed (optical and SEM) conditions and Steel 2 in (b) hot rolled (optical), (d) and (f) warm rolled and annealed (optical and SEM) conditions. (Online version in color.)

Table 4. The ferrite grain sizes of Steel 1 and Steel 2 after different processing steps.

	Steel 1		Steel 2
	Hot rolled	Cold rolled and annealed	Hot rolled	Warm rolled and annealed
Grain size (μm)	477 ± 175	39 ± 14	406 ± 187	33 ± 13

Table 5. Amounts of ferrite and Fe₃Al phases in Steel 1 and Steel 2 in different processing conditions.

Area %	Steel 1		Steel 2
	Hot rolled	Cold rolled and annealed	Hot rolled	warm rolled and annealed
ferrite	100	96.49 ± 0.62	97.59 ± 0.10	91.58 ± 0.44
Fe3Al	…	3.50 ± 0.62	2.38 ± 0.10	8.41 ± 0.51

Fig. 4.

SEM images and EDS profiles of (a) cold rolled and annealed sample of Steel 1, (b) warm rolled and annealed sample of Steel 2, (c) Ti rich precipitates along the grain boundary of cold rolled and annealed sample of Steel 1. (Online version in color.)

Fig. 5.

The XRD plots of Steel 1 and Steel 2 in different processing conditions. (Online version in color.)

3.3. Evolution of Tensile Properties and Hardness during Hot, Cold or Warm Rolling and Annealing

The room temperature engineering stress-strain curves (Figs. 6(a) and 6(b)), true stress-strain curves (Fig. 6(c)) and rate of strain hardening versus plastic strain curves (Fig. 6(d)) of Steel 1 and Steel 2 are given in Fig. 6, and the corresponding yield strength (YS), ultimate tensile strength (UTS), total elongation (TE) and yield ratio (YR) are given in Table 6. Hot rolled samples of Steel 1 and Steel 2 exhibit 0.51% and 1.7% elongations. Warm rolled and annealed sample of Steel 2 shows discontinuous yielding and TE is less than that of cold rolled and annealed sample of Steel 1 (Fig. 6(b)). The higher Al content in Steel 2 with greater area fraction of Fe₃Al intermetallic (Table 5) than in Steel 1 results in an increase in the strength and decrease in the elongation of the former. The maximum elongation of 17% is obtained in cold rolled and annealed sample of Steel 1. The strain hardening curves (Fig. 6(d)) show that the rate of strain hardening of the warm rolled and annealed sample of Steel 2 is slightly higher than that of cold rolled and annealed specimen of Steel 1. The microhardness (HV_0.1) and macrohardness (HV₁₀) values of Steel 1 and Steel 2 of hot rolled, cold/warm rolling and annealing are given in Table 7. In both the steels, the microhardness values after cold/warm rolling and annealing (900°C, 5 minutes) are similar to those in hot rolled condition. However, the micro- and macrohardness values of Steel 2 (hot rolled, warm rolled and annealed sample) are higher than Steel 1 (hot rolled, cold rolled and annealed sample. The yield ratio (YS/UTS) reflects the ability to harden the materials during plastic deformation i.e., decreasing yield ratio implies strengthening of the steels. After annealing of both steels the yield ratio decreases (Table 6) and hence also an improvement in their formability is expected.

Fig. 6.

(a) Engineering stress-strain curves of hot rolled (b) Engineering stress-strain curves of cold rolled/warm rolled and annealed (c) True stress-strain curves of cold rolled/warm rolled and annealed (d) Rate of strain hardening curves of Steel 1 and Steel 2 after cold /warm rolling and annealing at 900°C for 5 minutes. (Online version in color.)

Table 6. Variation of yield strength (YS), ultimate tensile strength (UTS) and total elongation of Steel 1 and Steel 2 with different processing conditions.

Steel name	Sample condition	Yield strength (MPa)	Ultimate tensile strength (MPa)	Total elongation (%)	Yield ratio (YS/UTS)
Steel 1	Hot rolled	310	328	0.51	0.94
	Cold rolled and annealed	337	444	17	0.76
Steel 2	Hot rolled	488	544	1.7	0.90
	Warm rolled and annealed	501	607	11.7	0.82

Table 7. Variation of Vickers microhardness at load 100 gf (HV_0.1) and macrohardness at load 10 kgf (HV₁₀) of Steel 1 and Steel 2 in different processing conditions.

	Steel 1				Steel 2
	Hot rolled		Cold rolled and annealed		Hot rolled		Warm rolled and annealed
	ferrite	Fe3Al	ferrite	Fe3Al	ferrite	Fe3Al	ferrite	Fe3Al
Micro hardness (HV0.1)	201 ± 11	204 ± 6	197 ± 5	213 ± 5	263 ± 2	270 ± 11	258 ± 12	267 ± 11
Macro hardness (HV10)	167 ± 4		161 ± 13		216 ± 7		204 ± 9

The SEM fracture surface images of tensile tested samples are shown in Figs. 7(a)–7(d). In Steel 1, in hot rolled condition, the cleavage mode fracture is dominant (Fig. 7(a)) as coarse and planar cleavage facets are visible and in cold rolled and annealed sample both cleavage mode fracture, and quasi cleavage mode fracture (wavy cleavage) and a few secondary cracks are observed (Fig. 7(b)). In Steel 2, in hot rolled stage, cleavage mode fracture and large secondary cracks are detected (Fig. 7(c)), whereas after warm rolling and annealing, cleavage mode, and quasi or wavy cleavage mode fracture is visible and modified secondary cracks are observed (Fig. 7(d)).

Fig. 7.

The SEM fracture images of (a) hot rolled, and (b) cold rolled and annealed samples of Steel 1; (c) hot rolled, and (d) warm rolled and annealed samples of Steel 2. (Online version in color.)

4. Discussion

4.1. Rolling and Annealing Strategies

The choice of deformation strategy strongly depends on the amount of Al contents of low density steels. In Steel 1 (6.8 mass% Al), sample showed no cracks even after 60% cold rolling (Fig. 8(a)). However, in Steel 2 (9.7 mass% Al), centre and edge cracks appeared after 69% of cold rolling (Fig. 8(b)) due to the presence of relatively higher amount of Fe₃Al intermetallic.^7,9,14) Moreover, both micro- and macrohardness of Steel 2 are higher than Steel 1 (Table 7) indicating higher solid solution hardening of ferrite by Al. Therefore, the cold rolling of the higher Al-containing Steel 2 results in cracking due to solid solution hardening of Al in ferrite. Consequently, Steel 2 was deformed through warm rolling (rolling temperature 250°C) to avoid prior crack formation. If the presence of Fe₃Al intermetallic precipitates is limited to only the ferrite grains (which was observed in our case), the problem of cracking in high Al containing low density Steel 2 could be minimised.¹⁸⁾ The best preventive method to stop the formation of cracks during cold rolling in high Al low density steels could be to control the lengthening and thickening (i.e. growth) of Fe₃Al intermetallic by fast cooling from the hot rolling temperature to coiling temperature through the Fe₃Al intermetallic formation temperature.¹⁶⁾ Just below Fe₃Al intermetallic formation temperature, the lengthening and thickening rates are higher. Intergranular Fe₃Al intermetallic forms in both steels after annealing. The strained grains after rolling, are replaced by new strain free grains during annealing above the recrystallization temperature. Rana et al.¹⁹⁾ found that in 6.57Al-3.34Mn-0.179C mass% alloy, the recrystallization of the ferrite phase complete at 850°C. However, recrystallization temperature of present steels could be somewhat higher. During solidification, partition ratios of Al in Steel 1 (1.080) and Steel 2 (1.058) are lowest compered to C and Ti given in Table 3, which indicate that Al remains in ferrite matrix.

Fig. 8.

Cold rolling behavior of (a) Steel 1 (60% reduction), (b) Steel 2 (69% reduction). (Online version in color.)

Figures 3(a) and 3(b) show the recrystallized large ferrite grains of Steel 1 and Steel 2 in hot rolled conditions. But, after cold/warm rolling and annealing (900°C, 5 minutes), the ferrite grains refine significantly in both the steels (Table 4). This grain refinement and the alloying elements Al and C also affected the tensile properties of Steel 1 and Steel 2. Cold and warm rolling were done below recrystallization temperatures, where strain hardening is not eliminated. Strain energy, present in the cold or warm rolled samples, can promote the nucleation of new strain free grains during annealing and refine the ferrite grains. Therefore, proper selection of rolling and annealing conditions for both steels based on Al and C contents can improve the tensile properties (Figs. 6(a) and 6(b)). The annealing temperature can change the shape and size of the Fe₃Al intermetallic as Fig. 3(f) illustrates for Steel 2 where the intergranular Fe₃Al intermetallic after warm rolling and annealing (900°C for 5 minutes) can be observed.

4.2. Microstructural Feature and Tensile Properties

The major phase in both the steels before annealing and after annealing are ferrite (Table 5). According to Hall-Petch relation (given later in Eq. (1)), reduction of ferrite grains after annealing improves the yield strength (Table 6) of Steel 1 and Steel 2. The total elongation of Steel 1 and Steel 2 increases from 0.51% (hot rolled Steel 1) and 1.7% (hot rolled Steel 2) to 17% (cold rolled and annealed Steel 1) and 11.7% (warm rolled and annealed Steel 2) respectively. The strain hardening in both steels after annealing (900°C, 5 minutes) gradually decreases due to activation of multiple slip systems (Fig. 6(d)). Similar trend has also been reported by Pramanik et al.¹⁴⁾ The rate of strain hardening of Steel 2 is higher than that of Steel 1 (Fig. 6(d)). This could be due to higher amount of Fe₃Al in Steel 2 than in Steel 1 (Table 5) implying more dislocation-precipitate interactions. After annealing of Steel 1 and Steel 2, approximately 19% and 8.8% reduction in yield ratios are measured. Therefore, the stretch-formability of Steel 1 after annealing could be better than that of Steel 2. The yield point behaviour was seen in Steel 2 after annealing (inset of Fig. 6(b)) and it could also deteriorate the formability of Steel 2. It could therefore, be argued that after rolling and annealing, the formability of both steels improves from hot rolled condition.

4.3. Strengthening Mechanisms

4.3.1. Grain Refinement

The average grain size of Steel 1 and Steel 2 after cold/warm rolling and annealing (900°C, 5 minutes) decreases approximately 92% (Table 4). In the cold rolled sample of Steel 1, the grains can be associated with higher strain energy and dislocations densities, than warm rolled sample of Steel 2, exhibiting greater driving force of nucleation of new grains in Steel 1 than in Steel 2 after annealing 900°C for 5 minutes. The strength and ductility of Steel 1 and Steel 2 increase due to grains refinement shown in Figs. 6(a) and 6(b). Dislocation pile up around the grain boundaries can also raise the strength of both steels. The grain boundary itself is the source of dislocations and it is the potential site for diffusion, and precipitation of the Fe₃Al intermetallic in rolled and annealed samples of Steel 1 and Steel 2 (Figs. 3(c) and 6(d)) improving the yield strength.

The Hall-Petch relation is given below:²⁰⁾   

$${\sigma }_{\mathrm{y}}={\sigma }_{\mathrm{i}}+\mathrm{K}{\mathrm{d}}^{-1/2}$$(1)

Where, σ_y is the yield stress, σ_i is friction stress, K is the strengthening coefficient and d is the average grain diameter. In the present study the friction stress σ_i = 100 MPa and K in the range of 550 to 600 MPa.μm^1/2 are considered for both steels (which have ferrite as the matrix) following the work of Takeda et al.²¹⁾ The yield strength of both the steels after rolling and annealing (900°C for 5 minutes) are calculated in the following:   

$${\sigma }_{\mathrm{y}}=100+590{\left(39\right)}^{-1/2}\left(\mathrm{M}\mathrm{P}\mathrm{a}\right),\mathrm{\hspace{0.17em} }=194.47\mathrm{\hspace{0.17em} }\mathrm{M}\mathrm{P}\mathrm{a}\mathrm{\hspace{0.17em} }\mathrm{f}\mathrm{o}\mathrm{r}\mathrm{\hspace{0.17em} }\mathrm{S}\mathrm{t}\mathrm{e}\mathrm{e}\mathrm{l}\mathrm{\hspace{0.17em} }1$$(2)

  

$${\sigma }_{\mathrm{y}}=100+590{\left(33\right)}^{-1/2}\left(\mathrm{M}\mathrm{P}\mathrm{a}\right),\mathrm{\hspace{0.17em} }=202.70\mathrm{\hspace{0.17em} }\mathrm{M}\mathrm{P}\mathrm{a}\mathrm{\hspace{0.17em} }\mathrm{f}\mathrm{o}\mathrm{r}\mathrm{\hspace{0.17em} }\mathrm{S}\mathrm{t}\mathrm{e}\mathrm{e}\mathrm{l}\mathrm{\hspace{0.17em} }2$$(3)

In Eqs. (2) and (3), the grain sizes in cold/warm rolled and annealed conditions have been taken from Table 4. The grain boundary strengthening according to Eq. (1) is 194.47 MPa for Steel 1 and 202.70 MPa for Steel 2. The predicted yield stresses in Steels 1 and 2 are lower than the measured values (Table 6) due to presence of other strengthening mechanisms.²²⁾ The higher grain refinement is responsible for higher yield strength of Steel 2 than Steel 1 suggesting the importance of the microstructural features than the steel chemistry in the present study.

4.3.2. Solid Solution Strengthening

To acquire better yield strength after rolling and annealing, solute solution strengthening due to Al atoms offers important contribution.²³⁾ The effect of solid solution strengthening due to Al atoms can be calculated by.²⁴⁾   

$$\mathrm{\Delta }{\sigma }_{\mathrm{s}\mathrm{s}}={\mathrm{k}}_{\mathrm{s}\mathrm{s}}{\mathrm{C}}_{\mathrm{A}\mathrm{l}}$$(4)

Where, k_ss (MPa/at.%) is the solid solution strengthening coefficient (it combined the effect of both modulus misfit and size misfit), C_Al is the concentration of the Al (at.%) solute atoms (Steel 1 and Steel 2: 13.12 at.% and 18.18 at.%). For the calculation of the yield strength contribution (Δσ_ss) value for Steel 1 and Steel 2, solid solution strengthening coefficient k_ss, is taken 9 MPa/at.%.²⁴⁾ According to Eq. (4), the solid solution strengthening effect for Steel 1 and Steel 2 are 118.17 MPa and 163.62 MPa, explaining the higher strength of Steel 2 than Steel 1 (Fig. 6(b)).

4.3.3. Strain Hardening

The true stress-strain curve (Fig. 6(c)) and the rate of strain hardening-plastic strain curve (Fig. 6(d)) describe the strain hardening behaviour of the both steels. Figure 6(d) shows that Steel 2 exhibits higher (due to higher precipitate-dislocation interaction) rate of strain hardening than Steel 1 due to higher contribution of precipitate interaction in the former. The yield point phenomena observed in Steel 2 after annealing, shown in inset (Fig. 6(b)). After annealing, the higher rate of strain hardening in Steel 2 presumably is due to the presence of higher amount of Fe₃Al intermetallic than Steel 1 (Table 5) and higher solid solution strengthening owing to greater amount of carbon and aluminium. This can hinder the motion of dislocations and cause accumulation of more dislocations (at non-shearable precipitates)^25,26) resulting in higher strength and lower total elongation of Steel 2 (warm rolled and annealing) than Steel 1 (cold rolled and annealing). Therefore, the formability of Steel 2 is expected to be lower than Steel 1. This also is supported by the work of Rana et al.¹⁸⁾ on duplex low-density steel.

Figure 9(a) and Table 8 show the amount of increase in the strength by different strengthening mechanisms of rolled and annealed samples of both steels. The strength increment is primarily due to grain refinement followed by solute solution strengthening. The predicted strengths of Steel 1 and Steel 2 is lower than the measured value due to the fact that dislocation-dislocation interaction and precipitation hardening are not taken into consideration for the calculation.^8,22) The characterisation of the microstructures in presence of chemical composition gradient at different stages of processing can allow accurate quantification of the different strengthening contributions removing the discrepancy between measured and the predicted strengths of the steels.

Fig. 9.

(a) Strengthening calculation after rolling and annealing of Steel 1 and Steel 2, (b) The yield strength verses total elongation plot of Steel 1 and Steel 2 and similar steels from literature for comparison purpose. (Online version in color.)

Table 8. The predicted strength contributions due to grain boundary strengthening, and solid solution strengthening in Steels 1 and 2.

Specimen	Strength contribution (MPa)		Predicted yield strength (MPa)	Measured yield strength (MPa)
	Grain refinement	Solid solution strengthening
Cold rolled and annealed, Steel 1	194.47	118.17	312.64	337
Warm rolled and annealed, Steel 2	202.70	163.62	366.32	501

Figure 9(b) exhibits the variation of yield strength and the total elongation of Steel 1 and Steel 2 after cold/warm rolling and annealing. The yield strengths and the total elongations of other steels with comparable range of C and Al are included in the Fig. 9(b). The Steels 1 and 2 present better combination of yield strength and ductility after cold/warm rolling and annealing than the existing results in the published literature.^{9,27,28,29,30)} The presence of higher carbon and aluminium in Steel 2 than Steel 1 increases the yield strength primarily due to solid solution strengthening. It offers to improvement in the combination of yield strength and total elongation and thus formability.

5. Conclusions

In this study, the effect of cold and warm rolling and annealing on the microstructures and tensile properties of low-density steels containing varying amounts of C and Al is explored. The major conclusions are:

• The Thermo-Calc analysis allowed prediction of the solidification behavior and micro segregation due to solute partitioning in low density steels. The thermodynamic calculations predict that the room temperature equilibrium phases in the steels are primarily ferrite, which matches very well with the observed microstructure. In addition to the steels compositions, the development of the microstructure with temperature controls the density of the steels which was estimated as a function of temperature using Thermo-Calc and validated experimentally.

• In Steel 1, the average ferrite grain size obtained in hot rolled sample is about 477 ± 175 μm and after cold rolled and annealing, the average ferrite grain size is 39 ± 14 μm. Similarly, for Steel 2, the average ferrite grain size in hot rolled sample is 406 ± 187 μm and in warm rolled and annealed sample is 33 ± 13 μm. After rolling and annealing (900°C for 5 minutes), ferrite grain refinement resulted in the increment in strength and total elongation in both the steels.

• In Steel 1 after cold rolling and annealing of hot rolled samples, the yield strength is increased from 310 MPa to 337 MPa and ultimate tensile strength is increased from 328 MPa to 444 MPa and the total elongation is raised from 0.51% to 17%. The highest yield strength (501 MPa) and ultimate tensile strength (607 MPa) are obtained in warm rolled and annealed sample of Steel 2. Significant grain refinements after cold or warm rolling and annealing resulted in improvement in the combination of strength and ductility of the steels. The fracture changed from planar cleavage in hot rolled conditions to quasi or wavy cleavage in cold or warm rolled and annealed conditions in Steels 1 and 2. With increase in the C and Al contents of the steel, the greater solid solution strengthening increases the strength of the steel.

• The prediction of strengths using the different strengthening mechanisms clarified the importance of the chemical composition and the microstructures. The finer characterisations of microstructures in presence of microsegregation during different stages of processing can eliminate the difference between measured and the predicted values of strength.

• In order to get the best combination of tensile properties, it is recommended to refine the grains size of the hot rolled samples of the low density steels. Based on the amount of Al content, Fe₃Al intermetallic can form leading to prior cracking during rolling. The cold rolling for Steel 1 (Al 6.8 mass%) and warm rolling for Steel 2 (Al 9.7 mass%) and then annealing (900°C for 5 minutes) help to achieve desirable strength and ductility combination for automotive body parts applications in ferritic low Carbon low density steels.

Acknowledgments

The financial support of the work has been received from Science and Engineering Research Board, Department of Science and technology, Government of India under core research grant (file no.: CRG/2020/001511). One of the authors (AK) is thankful to Professor P. C. Chakraborti, Metallurgical and Material Engineering Department, Jadavpur University, Kolkata-700032, India, for useful discussion and for the provision of the research facilities at Metallurgical and Material Engineering Department, Jadavpur University, Kolkata-700032, India and Centre of Excellence in Phase Transformation and Product Characterisation, Jadavpur University, Kolkata-700032, India, for the Thermo-Calc facility.

Nomenclature

FRT: Finishing rolling temperature

RT: Rolling temperature

RD: Rolling direction

TD: Transverse direction

ND: Normal direction

BCC: Body centred cubic

FCC: Face centred cubic

YS: Yield strength (MPa)

UTS: Ultimate tensile strength (MPa)

TE: Total elongation (%)

YR: Yield ratio

HV: Vickers Pyramid Number

σ_y: Yield strength (MPa)

σ_i: Friction stress (MPa)

K: Hall-Petch strengthening coefficient (MPa.μm^−1/2)

d: Average grain diameter (μm)

Δσ_ss: Yield strength due to solid solution strengthening (MPa)

k_ss: Solid solution strengthening coefficient (MPa/at.%)

C_Al: Concentration of Aluminium (at.%)

References

• 1)   G.  Frommeyer and  U.  Brüx: Steel Res. Int., 77 (2006), 627. https://doi.org/10.1002/srin.200606440
• 2)   I.  Gutierrez-Urrutia: ISIJ Int., 61 (2021), 16. https://doi.org/10.2355/isijinternational.ISIJINT-2020-467
• 3)   N.  Lim,  H.  Park,  S.  Kim and  C.  Park: Met. Mater. Int., 18 (2012), 647. https://doi.org/10.1007/s12540-012-4012-0
• 4)   C. L.  Lin,  C. G.  Chao,  H. Y.  Bor and  T. F.  Liu: Mater. Trans., 51 (2010), 1084. https://doi.org/10.2320/matertrans.M2010013
• 5)   W. K.  Choo,  J. H.  Kim and  J. C.  Yoon: Acta Mater., 45 (1997), 4877. https://doi.org/10.1016/S1359-6454(97)00201-2
• 6)   H.  Huang,  D.  Gan and  P. W.  Kao: Scr. Metall. Mater., 30 (1994), 499. https://doi.org/10.1016/0956-716X(94)90610-6
• 7)   S.  Chen,  R.  Rana,  A.  Haldar and  R. K.  Ray: Prog. Mater. Sci., 89 (2017), 345. https://doi.org/10.1016/j.pmatsci.2017.05.002
• 8)   K.  Ishida,  H.  Ohtani,  N.  Satoh,  R.  Kainuma and  T.  Nishizawa: ISIJ Int., 30 (1990), 680. https://doi.org/10.2355/isijinternational.30.680
• 9)   R.  Rana,  C.  Liu and  R. K.  Ray: Scr. Mater., 68 (2013), 354. https://doi.org/10.1016/j.scriptamat.2012.10.004
• 10)   A.  Zargaran,  H. S.  Kim,  J. H.  Kwak and  N. J.  Kim: Scr. Mater., 89 (2014), 37. https://doi.org/10.1016/j.scriptamat.2014.06.018
• 11)   D. G.  Morris,  M. A.  Muñoz-Morris and  L. M.  Requejo: Mater. Sci. Eng. A, 460–461 (2007), 163. https://doi.org/10.1016/j.msea.2007.01.014
• 12)   J.  Liu,  Z.  Liu,  C.  Gao,  Z.  Wang,  X.  Liu and  G.  Wang: J. Mater. Process. Technol., 167 (2005), 132. https://doi.org/10.1016/j.jmatprotec.2004.10.010
• 13)   M.  Zhang,  H.  Xu,  W.  Cao and  H.  Dong: ISIJ Int., 56 (2016), 861. https://doi.org/10.2355/isijinternational.ISIJINT-2015-425
• 14)   S.  Pramanik and  S.  Suwas: J. Mater. Eng. Perform., 31 (2022), 7449. https://doi.org/10.1007/s11665-022-06792-7
• 15)   C.  Castan,  F.  Montheillet and  A.  Perlade: Scr. Mater., 68 (2013), 360. https://doi.org/10.1016/j.scriptamat.2012.07.037
• 16)   S. S.  Sohn,  B.  Lee,  S.  Lee and  J.  Kwak: Acta Mater., 61 (2013), 5626. https://doi.org/10.1016/j.actamat.2013.06.004
• 17)   S. S.  Sohn,  B.  Lee,  S.  Lee,  N. J.  Kim and  J.  Kwak: Acta Mater., 61 (2013), 5050. https://doi.org/10.1016/j.actamat.2013.04.038
• 18)   R.  Rana: High-Performance Ferrous Alloys, Springer, Cham, (2021), 211. https://doi.org/10.1007/978-3-030-53825-5
• 19)   R.  Rana,  C.  Liu and  R. K.  Ray: Philos. Mag. Lett., 94 (2014), 127. https://doi.org/10.1080/09500839.2013.878046
• 20)   B.  Fu,  C.  Pei,  H.  Pan,  Y.  Guo,  L.  Fu and  A.  Shan: Mater. Res. Express, 7 (2020), 116516. https://doi.org/10.1088/2053-1591/abc79f
• 21)   K.  Takeda,  N.  Nakada,  T.  Tsuchiyama and  S.  Takaki: ISIJ Int., 48 (2008), 1122. https://doi.org/10.2355/isijinternational.48.1122
• 22)   M. X.  Yang,  F. P.  Yuan,  Q. G.  Xie,  Y. D.  Wang,  E.  Ma and  X. L.  Wu: Acta Mater., 109 (2016), 213. https://doi.org/10.1016/j.actamat.2016.02.044
• 23)   S.  Pramanik,  S.  Koppoju,  A. V.  Anupama,  B.  Sahoo and  S.  Suwas: Mater. Sci. Eng. A, 712 (2018), 574. https://doi.org/10.1016/j.msea.2017.10.056
• 24)   Q.  Lu,  W.  Xu and  S.  van der Zwaag: Comput. Mater. Sci., 84 (2014), 198. https://doi.org/10.1016/j.commatsci.2013.12.009
• 25)   A.  Deschamps,  S.  Esmaeili,  W. J.  Poole and  M.  Militzer: J. Phys. IV France, 10 (2000), Pr6-151. https://doi.org/10.1051/jp4:2000626
• 26)   M.  Koyama,  T.  Sawaguchi and  K.  Tsuzaki: ISIJ Int., 58 (2018), 1383. https://doi.org/10.2355/isijinternational.ISIJINT-2018-237
• 27)   S. Y.  Han,  S. Y.  Shin,  S.  Lee,  N. J.  Kim,  J. H.  Kwak and  K. G.  Chin: Metall. Mater. Trans. A, 42 (2011), 138. https://doi.org/10.1007/s11661-010-0456-3
• 28)   R. G.  Baligidad: J. Mater. Sci., 39 (2004), 5599. https://doi.org/10.1023/B:JMSC.0000039298.29397.3a
• 29)   R. G.  Baligidad: Mater. Sci. Eng. A, 368 (2004), 131. https://doi.org/10.1016/j.msea.2003.10.243
• 30)   H.  Song,  Y.  Kwon,  S. S.  Sohn,  M.  Koo,  N. J.  Kim,  B.-J.  Lee and  S.  Lee: Mater. Sci. Eng. A, 730 (2018), 177. https://doi.org/10.1016/j.msea.2018.06.003