Dynamic Recrystallization and Grain Refinement of ...

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Dynamic Recrystallization and Grain Refinement of Fe-P-C-Si and Fe-P-C-Si-N Steels Yashwant Mehta, S.K. Rajput, G.P. Chaudhari, and Vikram V. Dabhade (Submitted April 26, 2017; in revised form July 3, 2018) Grain refinement is an effective technique to improve the mechanical properties of steels. In the present work, single-pass hot compression experiments have been conducted on two different compositions of high phosphorous steels to study the microstructural evolution and ferrite grain refinement at various strain rates and deformation temperatures, i.e., 0.01-10 s21 and 750-1050 °C, respectively. Optical metallography has been employed to understand the physical processes that take place during hot deformation process. The results indicate that when these compositions of high phosphorous steels are worked at relatively low temperatures in the intercritical regions, a ferrite grain size of 5-7 lm is obtained. It is observed that the grain size decreases with an increase in strain rate and with the decrease in deformation temperature. Based on the values of stress exponent (n) obtained in the present work, dislocation creep is identified as the deformation mechanism. The activation energies for deformation of these two types of high phosphorous steels have been calculated. The effect of the alloying elements on the stress–strain curve, microstructure, and grain refinement has been discussed. Keywords

activation energy, dynamic recrystallization, grain refinement, high phosphorous steel, thermomechanical processing

1. Introduction Steels with optimized mechanical properties can be produced using thermomechanical processing. Usually, controlled rolling is used to produce steels with improved yield strength at low costs (Ref 1). However, fine-grained structure can be produced using three mechanisms: (a) strain-induced transformation (SIT), (b) transformation from dynamically recrystallized austenite, and (c) dynamic recrystallization of ferrite (Ref 2). Several studies have confirmed the achievement of fine ferrite via strain-induced transformation in plain carbon steels (Ref 3-6) and microalloyed steels (Ref 7, 8). These researchers have hot-worked the steel either just above the Ar3 temperature or in the intercritical region to obtain fine grain size. In order to produce fine-grained steel, it is important to understand the high-temperature deformation behavior of the material. Saadatkia et al. (Ref 9), while studying the hot deformation behavior of low and medium carbon steels identified three types of dynamic recrystallization behaviors, namely single peak, multiple transient steady state, and multiple peak. Zhao et al. (Ref 10) evaluated the hot deformation behavior of vanadium microalloyed steels and found dislocation climb as the ratecontrolling mechanism at low strain rates and cross-slip at high strain rates. Mirzadeh et al. (Ref 11) studied the hot deformation behavior of a medium carbon microalloyed steel and Yashwant Mehta, G.P. Chaudhari, and Vikram V. Dabhade, Department of Metallurgical and Materials Engineering, Indian Institute of Technology, Roorkee 247667, India; S.K. Rajput, Department of Mechanical Engineering, Bundelkhand Institute of Engineering and Technology, Jhansi 284128, India. Contact e-mail: [email protected].

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reported 394 kJ/mol as the deformation activation energy. Other researchers (Ref 12, 13) used various constitutive modeling methods to predict flow stress for medium carbon microalloyed steels during hot deformation. Phosphorous has a marked solid solution strengthening effect in ferrite which is of the same order as the interstitial elements, carbon and nitrogen (Ref 14). Phosphorous improves corrosion resistance by forming a passive surface film as is found in the Delhi iron pillar (Ref 15, 16). Grain refinement has been reported (Ref 17) to be very effective in enhancing toughness of ferritic steels. Accordingly, a suitable thermomechanical processing window needs to be found in the domain of strain rate and deformation temperature to achieve a finegrained material. Although hot compression behavior of high phosphorous steels has been studied by Kim et al. (Ref 18), very little or no work has been reported that estimates the combined effect of strain rate and deformation temperature on the development of fine ferrite grain size in high phosphorous steels. In view of this, hot compression tests were conducted on two types of low carbon high phosphorous steels to study the effect of different combinations of strain rates and deformation temperatures on the refinement of ferrite grains. Effect of the alloying elements on the stress–strain curves, microstructures and grain refinement of high phosphorous steels is also evaluated in the present study.

2. Experimental Procedure 2.1 Alloy Manufacturing An induction furnace of 300 kg capacity was used to produce two different types of steels having high phosphorous content. Appropriate quantities of iron scrap, ferrosilicon, graphite, ferro-phosphorous, etc., were charged into the furnace for melting. Before casting, the desired amount of aluminum shots were added to the melt for deoxidation purposes. The melt so produced was cast and labeled as S1 steel. Once S1

of 50%, i.e., a true strain of 0.69. Finally, the hot-compressed specimens were water-quenched to freeze the developed microstructure and pre-empt any meta-dynamic or static process occurrences that might follow hot compression. The hot compression test was also performed at 1050 °C.

steel melt was cast, another steel melt was prepared to get S2type high phosphorous steel by adding about 0.1% Si, 0.3% Mn, and 0.011% N to the remaining melt in the furnace. Nitrogen was added to the melt in the form of nitrided manganese. The presence of silicon and nitrogen in the steel is expected to displace phosphorous away from the grain boundaries during its heat treatment in the intercritical region (Ref 19, 20), thereby reducing the segregation of phosphorous at the grain boundaries. Castings of simple geometrical dimensions as 400 mm 9 200 mm 9 40 mm were obtained. MAGELLAN optical emission spectrometer was used to determine the chemical composition of the two types of steels produced, and the results are given in Table 1. Each of S1 and S2 steel castings was cut into smaller pieces of 40 mm 9 30 mm 9 30 mm dimensions. Subsequently, these smaller pieces of S1 and S2 steel casting were soaked for half an hour at 1150 °C and hot-forged to make rods of about 12 mm diameter each. Appropriate-sized specimens to conduct the hot compression tests were machined from these forged rods.

2.4 Metallography The hot-compressed specimens were sliced at the center along the compressive axis for making specimens for metallographic analysis. These were first ground using belt grinder and then hand-polished using silicon carbide papers of grades 320, 800, 1200, 1500, and 2000. These specimens were subsequently velvet cloth polished by using 1-lm fine alumina slurry suspension. In order to reveal the microstructure, the polished specimens were etched with 2% nital solution. A light optical microscope (Leica DMI5000 M) was used to view the microstructure of the etched specimens. The grain size of the specimens was determined by the ASTM line intercept method (Ref 21).

2.2 Phase Transformation In order to establish the phase transformation temperatures for the high phosphorous steels produced and used in the present study, dilatometric testing was performed on GleebleÒ 3800 thermomechanical simulator. The tests were performed on cylindrical specimens of 85 mm length and 10 mm diameter. A K-type thermocouple was spot-welded at the center of each specimen to measure the temperature. The cylindrical specimens were heated to 1050 °C at a heating rate of 5 °C s1, soaked at this temperature for 10 s, and subsequently cooled to 25 °C at a cooling rate of 1 °C s1. The change in the volume of the specimen during cooling from 1050 to 25 °C was recorded in the form of dilatometric curves (Fig. 1) to establish the phase transformation temperatures for S1 and S2 steels.

3. Results and Discussions 3.1 Phase Transformation Temperatures The phase transformation temperatures Ar3 and Ar1 for the two types of high phosphorous steels S1 and S2 were established on the basis of their respective dilatometric curves obtained during the dilatometry studies conducted in a thermomechanical simulator. The first and second deviation points from linearity in the respective dilatometric curves of S1

2.3 Hot Compression Test

A r1

160 Ò

140

A thermomechanical simulator (Gleeble 3800) was used to conduct hot compression tests on S1 and S2 steel specimens. A vacuum of about 1 Pa pressure was maintained in the simulator, and the cylindrical specimens having dimensions of 15 mm length and 10 mm diameter were used during hot compression tests. A K-type thermocouple attached to the test specimen was used to measure and control the temperature while the experiment was being performed. A nickel-based lubricant and a graphite foil were used in between the test specimen and the anvil to decease the temperature gradient and friction. All the test specimens were heated to the austenitizing temperature of 1050 °C at a heating rate of 5 °C s1, and these were held at that temperature for 10 s and subsequently cooled at the rate of 1 °C s1 to the corresponding compression test temperatures of 1000, 950, 900, 850, and 750 °C, respectively. At each temperature level, the specimens were compressed at constant strain rates of 0.01, 0.1, 1.0, and 10 s1 for a total deformation

Dilation (μm)

120 100 80

Ar3

60 40 20 0 -20 0

500

1000

Temperature (°C) Fig. 1 Dilatometric curve for S1-type high phosphorous steel

Table 1 Chemical composition of the two types of steels produced and used in wt.% Steel type S1 S2

P

C

Si

Mn

Cr

Al

Cu

N

S

Fe

0.13 0.13

0.05 0.05

0.39 0.48

0.2 0.46

0.14 0.18

0.023 0.033

0.024 0.026

0.004 0.015

0.015 0.014

Balance Balance

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3.2 Stress–Strain Curves

and S2 steels correspond to their phase transformation temperatures. The values of Ar3 and Ar1 so established for the two steels S1 and S2 are given in Table 2. It is observed that for S2 high phosphorous steel, the Ar1 and Ar3 temperatures are lower than those of S1 high phosphorous steel. Nitrogen being an austenite stabilizer expands the gamma phase field in S2 steel and thereby causes its phase transformation temperatures to fall to lower values in comparison with S1 steel (Ref 22). The dilatometric curve for S1-type high phosphorous steel is shown in Fig. 1, and that for S2 steel is shown elsewhere (Ref 23).

The true flow stress of a material depends on the strain rate and temperature at which the hot compression test is carried out. Figure 2 shows the stress–strain flow curves for S1 steel obtained at different strain rates and different deformation temperatures. It is observed that the stress–strain curves obtained at strain rates of 0.01, 0.1, and 1 s1 and at the deformation temperatures of 800-1050 °C display a peak stress after work hardening. However, at the strain rate of 10 s1 the flow stress increases gradually up to the maximum true strain of 0.69 (Fig. 2a). As the strain rate increases, the work hardening rate rises and thereby makes the restoration process difficult. The increase in strain rate results in the rise of peak strain, i.e., the peak stress shifts toward higher strain while the deformation temperature is maintained constant. Higher deformation temperature favors easier movement of vacancies and dislocations, resulting in dynamic recrystallization (DRX). It is seen in Fig. 2 that at strain rates of 10, 1, 0.1, and 0.01 s1 the flow curve at 1050 °C shows DRX. The initial stages of deformation are related to the work hardening due to the generation and multiplication of dislocations taking place during hot compres-

Table 2 Phase transformation temperatures as determined from dilatometry studies of the two high phosphorous steels Phase transformation temperature

S1 steel

S2 steel

813 941

805 925

Ar1, °C Ar3, °C

240

950°C 800°C

Stress (MPa)

200 180 160

850°C

140 120 1050°C 950°C 900°C 850°C 800°C

100 80 60 0.0

0.1

0.2

0.3

0.4

0.5

900°C 800°C

120

1050°C

100

850°C

80

1050°C 950°C 900°C 850°C 800°C

60 40 0.0

0.6

Strain

(a)

0.1

0.2

0.3

0.4

0.5

–1

0.01 s

0.1 s

950°C

70

950°C

100

900°C

80

1050°C

70

850°C

60

1050°C 950°C 900°C 850°C 800°C

50 40

(c)

0.1

0.2

0.3

0.4

Strain

0.5

Stress (MPa)

Stress (MPa)

800 °C 90

30 0.0

0.6

Strain

(b)

–1

110

950°C

140

900°C 1050°C

Stress (MPa)

220

–1

160 1 s

–1

10 s

800°C

60

1050°C 900°C 850°C

50

1050°C 950°C 900°C 850°C 800°C

40

30 0.0

0.6

(d)

0.1

0.2

0.3

0.4

0.5

0.6

Strain

Fig. 2 Flow curves for S1 steel obtained at different deformation temperatures (800-1050 °C) and at strain rates: (a) 10, (b) 1, (c) 0.1, and (d) 0.01 s1

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sion. The rate of work hardening is greater than the rate of work softening up to the peak stress. Subsequently, softening processes due to dynamic recrystallization become dominant over work hardening. Therefore, the flow stress falls from the peak stress values to a steady state value. The formation of dislocation free grains through nucleation and growth during DRX is facilitated by higher deformation temperatures. Two peaks can be observed in the flow curve that has been obtained at a compression temperature of 1050 °C and at a strain rate of 0.01 s1 (Fig. 2d). Similar observations have been made by previous workers (Ref 24). Multiple peaks in the flow curve may correspond either to grain coarsening (Ref 25) or to the cases where the ratio of the initial to stable grain sizes is less than 2 (Ref 26, 27). At the deformation temperature of 950 °C, the flow curve shows work hardening taking place at the strain rate of 10 s1 (Fig. 2a), whereas DRX behavior is observed in the S1 steel as the strain rate decreases from 10 to 0.01 s1. Further, on reducing the deformation temperature to 900 °C, the level of the peak stress falls below the peak stress registered by the material at 950 °C. Furthermore, the flow curve shows work hardening at a strain rate of 10 s1, and DRX at strain rates of

1, 0.1, and 0.01 s1. At 850 °C, the flow curves obtained at 10 s1 (Fig. 2a) do not show work hardening beyond a strain of about 0.35, and the flow stress remains more or less constant. This type of non-work hardening behavior of the material is ascribed to dynamic recovery (DRV) of ferrite (Ref 28). Moreover, the flow curves show DRX at a compression temperature of 850 °C and at strain rates of 1, 0.1, and 0.01 s1. At the deformation temperature of 800 °C, the flow curves exhibit DRV at strain rates of 10, and 1 s1 and DRX at 0.1 and 0.01 s1. In general, the peak stress increases at 900950 °C due to work hardening and decreases at 800-850 °C due to the occurrence of DRV. However, the peak stress is low at 1050 °C for all the strain rates due to the occurrence of DRX. The flow curves for S2 steel acquired at different strain rates and at different deformation temperatures are presented in Fig. 3. It is observed in S2 steel that higher peak stresses correspond to (1) the strain rates of 10, 1, 0.1, and 0.01 s1 at deformation temperatures of 800 and 1050 °C, (2) the strain rates of 10 and 1 s1 at the deformation temperature of 950 °C, and (3) the strain rate of 10 s1 at the deformation temperature of 850 °C in comparison with the Fe-0.13P-0.05C-0.26Si0.2Mn steel studied earlier (Ref 29).

180

–1

200

160

900°C

180

900°C 850°C

140

950°C

160

1050°C

140 120

1050°C 950°C 900°C 850°C 800°C

100 80 60 0.0

0.1

0.2

(a)

0.3

0.4

0.5

100

850°C 1050°C 950°C 900°C 850°C 800°C

40 0.0

0.1

0.2

(b)

0.3

0.4

0.5

0.6

Strain 68

104

0.01 s–1

64

950°C

96

950°C

60

800°C

88 80

900°C

1050°C

72

850°C

64

1050°C 950°C 900°C 850°C 800°C

56 48 40 32 0.1

0.2

0.3

Strain

0.4

0.5

Stress (MPa)

Stress (MPa)

1050°C

60

0.6

Strain

0.0

800°C

120

80

112 0.1 s–1

(c)

–1

1s

950°C 800°C

Stress (MPa)

Stress (MPa)

220 10 s

56

800°C

52

1050°C

48 44

850°C 900°C

1050°C 950°C 900°C 850°C 800°C

40 36 32 0.0

0.6

(d)

0.1

0.2

0.3

0.4

0.5

0.6

Strain

Fig. 3 Flow curves for S2 steel acquired at different deformation temperatures (800-1050 °C) and at strain rates: (a) 10, (b) 1, (c) 0.1, and (d) 0.01 s1

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Figure 4 gives the comparison of flow curves of S1 and S2 steels at different deformation temperatures and strain rates. In comparison with S1 steel, lower peak stresses are observed in S2 steel at (1) the intercritical temperature of 900 °C and strain rates of 10, 0.1, and 0.01 s1 (Fig. 4a and c) and (2) the intercritical temperature of 850 °C and strain rate of 1 s1. However, S2 steel shows higher peak stress at 850 °C and strain rates of 10, 0.1, and 0.01 s1 (Fig. 4b and d). The observed difference in the values of peak stresses could be due to the presence of higher proportion of austenite in the microstructures in S2 steel developed at 850 °C and 10, 0.1 and 0.01 s1 strain rates. The higher proportion of austenite in the microstructure of S2 steel is attributed to the presence of higher amounts of nitrogen which acts as austenite stabilizer (Ref 22). Austenite is harder than ferrite at higher temperatures (Ref 30), and this leads to the higher peak stress values evident in S2 steel at the deformation temperature of 850 °C and strain rates of 10, 0.1, and 0.01 s1.

3.3 Microstructural Evolution Figure 5 presents the ferrite grain refinement and development of the microstructure of S1 steel compressed at different

temperatures and strain rates. At the deformation temperature of 1050 °C and strain rates of 0.1 and 0.01 s1, the microstructure developed is a mixture of polygonal and acicular ferrite. However, the microstructure obtained at the temperature of 1050 °C and strain rates of 1 and 10 s1 shows polygonal ferrite. Figure 2(a) and (b) confirms that the austenite has experienced DRX under these conditions. The microstructure at the deformation temperature of 950 °C and at the strain rate of 10 s1 shows equiaxed ferrite. The corresponding flow curve of S1 steel shows work hardening and recovery (Fig. 2a). This implies that ferrite is formed due to the transformation of workhardened or deformed austenite. The size of ferrite grains is smaller as compared to that obtained at the temperature of 1050 °C. The presence of some coarse ferrite grains in the microstructure at lower strain rates suggests that grain growth could have occurred, although nucleation rates are high at such high temperatures (Ref 8). The microstructures obtained at the deformation temperature of 950 °C and at strain rates of 1, 0.1, and 0.01 s1 show the presence of equiaxed ferrite due to DRX, which is also confirmed from the flow curves in Fig. 2. At the deformation temperature of 900 °C, finer ferrite grains are obtained as compared to those formed at 950 °C. A small amount of deformed and elongated proeutectoid ferrite

75

95 90

70

80

Stress (MPa)

Stress (MPa)

85

75 70 65 60 55

–1

0.1s 900-S1 900-S2

50 45 0.0

0.1

0.2

0.3

0.4

0.5

60 55 50 –1

45

0.1s

40

850-S1 850-S2

0.0

0.6

Strain

(a)

65

0.1

0.2

(b)

0.3

0.4

0.5

0.6

Strain 52

58 50

54

Stress (MPa)

Stress (MPa)

56

52 50 48

44

–1

–1

0.01s

0.01s 900-S1 900-S2

44

(c)

46

42

46

42 0.0

48

0.1

0.2

0.3

0.4

Strain

0.5

40 38 0.0

0.6

(d)

850-S1 850-S2 0.1

0.2

0.3

0.4

0.5

0.6

Strain

Fig. 4 The comparison of stress–strain flow curves of S1 and S2 steels at: (a) 900 °C and 0.1 s1, (b) 850 °C and 0.1 s1, (c) 900 °C and 0.01 s1, and (d) 850 °C and 0.01 s1

Journal of Materials Engineering and Performance

1050°C

950°C

900°C

850°C

800°C

0.01s−1

0.1 s−1

1s−1

10 s−1

S1

20μm Fig. 5 steel

Ferrite grain refinement and evolution of the microstructure at different compression strain rates and deformation temperatures in S1

grains are also observed in the microstructure. The flow curves obtained at 900 °C show different combinations of hardening, recovery, and recrystallization as the strain rate decreases (Fig. 2). This implies that the ferrite grains are formed from the deformed austenite grains at strain rates of 10 and 1 s1. The microstructures at deformation temperature of 900 °C and at strain rates of 0.1 and 0.01 s1 are transformed from recrystallized austenite grains, and the latter experienced grain growth due to the low strain rate and increased time that was available for growth (Ref 31). The resulting grain sizes are smaller than those obtained at higher temperatures due to the larger amount of undercooling, and hence a higher nucleation rate (Ref 5). At 850 °C, the microstructures possess equiaxed ferrite and deformed proeutectoid ferrite. The values of peak stress are lower than those obtained at 900 °C, thus pointing toward softening due to DRV at 10 s1, and DRX at 1, 0.1, and

0.01 s1. The flow curve (Fig. 2a) at 10 s1 shows a constant flow stress after a strain of 0.35, whereas the flow curves (Fig. 2b, c, and d) at 1, 0.1, and 0.01 s1, respectively, show dynamic recrystallization. Flow curves at 0.01 s1 show a constant flow stress after a strain of about 0.55, and the microstructure reflects equiaxed ferrite with some large grains that signify grain growth. Deformation in an undercooled condition produces a deformed and inhomogeneous ferrite since the proeutectoid ferrite cannot recrystallize easily. A few small dark colored grains are visible in the microstructure obtained at strain rate of 0.01 s1 which could not be resolved optically. However, these were later confirmed as pearlite grains by electron microscopy (Fig. 6). The deformation temperature of 850 °C falls in the intercritical phase field for S1 steel. The phase transformation temperatures shift to higher values if deformation accompanies cooling as is observed in AISI 1016

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steel and has been discussed elsewhere (Ref 28). The aforesaid elevation of transformation temperatures can be explained by either or all of the following reasons (Ref 28):

10 μm Fig. 6 SEM micrograph showing pearlite grains in S1 steel at 850 °C and 0.01 s1

1050°C

950°C

The presence of pearlite in the microstructure developed at 850 °C can be explained in the light of the above facts. At 800 °C, the microstructure consists of subgrains, deformed proeutectoid ferrite, and fine equiaxed ferrite. The

900°C

850°C

800°C

0.01s−1

0.1 s−1

1s−1

10 s−1

S2

1. The Gibbs free energy of the austenite phase increases due to plastic deformation. 2. The presence of high dislocation density increases the number of sites for heterogeneous nucleation and diffusion of carbon. 3. The austenite grains are refined by dynamic recrystallization.

20μm Fig. 7

Evolution of the microstructure and ferrite grain refinement at diverse strain rates and hot compression temperatures of S2 steel

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16

16

S1

14

12 10 8

10 s–1 –1 1s –1 0.1 s –1 0.01 s

6

Ferrite grain size (μm)

Ferrite grain size (μm)

14

4

12 10 8

–1

10 s –1 1s –1 0.1 s –1 0.01 s

6 4

800

850

900

950

1000

1050

800

Deformation temperature (°C)

(a) Fig. 8

S2

850

900

950

1000

1050

Deformation temperature (°C)

(b)

Effects of various strain rates and hot compression temperatures on the ferrite grain size of (a) S1 and (b) S2 steels 180 160

Stress (MPa)

140 120 100 80 60 40 20 0 0.0

(a) Fig. 9

750°C, 1 s

–1

0.1

(b)

0.2

0.3

0.4

0.5

0.6

Strain

S1 steel specimen compressed at 750 °C and 1 s1 (a) Microstructure (b) flow curve

corresponding flow curves (Fig. 2a and b) indicate DRV at 10 and 1 s1; and DRX at 0.1 and 0.01 s1 (Fig. 2c and d). The amount of pearlite grains in the microstructure have slightly increased due to the increased diffusion of carbon owing to the increased time spent in cooling from 1050 to 800 °C. Figure 7 shows the evolution of the microstructure and ferrite grain refinement at diverse strain rates and hot compression temperatures in S2 steel. The microstructures obtained are similar to those developed in the case of S1 steel. The effect of deformation temperature on the ferrite grain size of S1 and S2 steels is shown in Fig. 8(a) and (b), respectively. The largest ferrite grains are obtained at 1050 °C and the smallest at 800 °C. However, at 800 °C the largest ferrite grains are obtained at a strain rate of 0.01 s1. This is due to the availability of greater amount of time for grain growth at that strain rate. Mostly, the ferrite grains obtained at 0.1 s1 are found to be smaller than those obtained at 0.01 s1 in both S1 and S2 steels. The grain sizes obtained at 1 s1 and at different hot compression temperatures are intermediate between the grain sizes achieved at 0.1 and 0.01 s1. This could be due to the adiabatic rise in temperature. In general, the increase in strain rate from 0.01 to 0.1 s1 causes a decrease in the ferrite grain sizes in both the steels. While comparing the grain sizes of ferrite developed at different strain rates, it is

found that the grain size of S1 steel is larger than that of S2 steel at the strain rate of 0.01 s1. It implies that the increased alloying (addition) of Si, N, and Mn in S2 steel causes a decrease in grain size at 0.01 s1 strain rate and at the deformation temperatures of 800, 850, 900, 950, and 1050 °C. However, the trend is not consistent at other strain rates because of the adiabatic rise in temperature. The smallest ferrite grain size of 5.5 lm is observed in S1 steel at the strain rate of 1 s1 and deformation temperature of 800 °C, whereas in S2 steel the smallest ferrite grain size of 5.7 lm is obtained at a strain rate of 0.1 s1 and deformation temperature of 800 °C. The experiments performed on S1 and S2 steels at the hot compression temperature of 750 °C in order to evaluate the effect of strain rates of 0.1 and 1 s1 reveal that the size of the ferrite grains developed in S1 steel is the smallest, having the average grain size of about 3.7 lm. The microstructure developed and the corresponding flow curve obtained for S1 steel are presented in Fig. 9. The mixed type microstructure of S1 steel resulting from hot compression at 850 °C using the strain rate of 10 s1 can be described on the basis of ASTM E1181 standard (Ref 32) as: duplex, wide range, 55% ASTM No. 11.5, 45% ASTM No. 6.5. Further, a few pearlite grains are observed in some of the microstructures developed at deformation temperatures of 800 and 850 °C. The percentage of field

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Table 3 Adiabatic rise in temperature during compression of S1 and S2 steels dT, °C at the strain rate of S1 steel Deformation temperature, °C 750 800 850 900 950 1000 1050

S2 steel

0.01 s21

0.1 s21

1 s21

0.01 s21

0.1 s21

1 s21

3.8 6.9 4.2 4.2 0.6 2 1.4

5.7 5.3 4.6 1.9 4.5 4.5 4.2

13.2 17.7 18.3 19.2 22.1 11.7 12.1

3.7 5.2 2.6 3.7 1.1 0.9 1.4

8.6 3.5 6.3 2.9 4.3 4.2 3.9

27 19.2 17.7 22.6 19.4 18.4 10.2

Table 4 Peak stress corrected for adiabatic rise in temperature for S1 and S2 steels Peak stress corrected for adiabatic rise in temperature S1

S2

Temp., °C

1 s21

0.1 s21

0.01 s21

1 s21

0.1 s21

0.01 s21

750 800 850 900 950 1000 1050

164.68 136.11 109.18 123.03 151.92 140.05 122.85

130.67 95.01 70.10 88.67 110.87 98.08 84.26

92.57 66.29 48.57 58.79 75.48 64.45 56.19

208.51 139.91 104.16 139.58 158.19 141.04 121.99

127.48 95.35 73.11 86.62 107.96 100.68 85.14

78.99 62.16 50.73 51.54 66.08 65.36 55.92

area occupied by the pearlite phase in S1 steel hot-compressed at 0.01 s1 and 850 °C when determined according to the ASTM E562 is found to be mere 7.5% (Ref 33) and as such the pearlite grains were not taken into account while measuring and comparing the overall grain sizes.

3.4 Activation Energy for Hot Deformation The activation energy for hot compression indicates the problems involved in the atomic relocations in mechanisms like recrystallization (Ref 18, 34, 35). The universal constitutive equation proposed by Sellars and McTegart (Ref 36) is used to examine the influence of higher alloying on activation energy in S2 steel. The proposed equation can be written as:   Qd n _ ðEq 1Þ Z ¼ A½sinhðarp Þ ¼ e exp RT In this Eq 1, Z is the Zener–Hollomon parameter, Qd is the activation energy for deformation, R is the universal gas constant, T is the absolute temperature, e_ is the strain rate, and rp is the peak stress at the specified strain rate and temperature. A (s1), a (MPa1), and n are the material constants where n is the stress exponent, and 0.012 is taken as the value of a as obtained for comparable steels from the sinh type equation (Ref 37). The activation energy is calculated using the Arrhenius graph assuming that microstructures do not vary. The adiabatic rise in temperatures is tabulated in Table 3. For S1 steel, the maximum rise in temperature is 22.1 °C at 950 °C and 1 s1. However, the maximum rise in temperature for S2 steel is 27 °C at 750 °C and 1 s1. The peak stresses were corrected

Journal of Materials Engineering and Performance

for adiabatic rise in temperature during hot compression, and the corrected values are tabulated in Table 4. It is seen that the peak stresses decrease as temperatures increases from 750 to 850 °C since the microstructures are predominantly composed of ferrite phase and ferrite softens at higher temperatures. The amount of austenite phase fraction starts increasing at the expense of ferrite phase fraction between 900 and 950 °C. Since austenite is harder than ferrite at higher temperatures, the peak stress values increase. Subsequently, at higher temperatures austenite softens and the peak stress values start decreasing up to 1050 °C. The corrected peak stresses are correlated with the reciprocal of the absolute temperature for a large number of combinations of temperatures and strain rates at which hot deformation is performed for a true strain of 0.69. The corrected peak compressive stresses at strain rates of 1, 0.1, and 0.01 s1 and temperatures of 750, 800, 850, 900, 950, 1000, and 1050 °C are determined, and the parameters and constants are calculated. The calculated parameters and constants are represented in Table 5. The slope of ln (sinh(ar)) versus 10,000/T for the temperature range 850-950 °C is found to be negative for both the steels under consideration. High-temperature ferrite is known to have lower hardness than austenite (Ref 30). This gives rise to very low activation energies of 37.6 and 50.5 kJ/mol for S1 steel and S2 steel, respectively, for overall range of 750-1050 °C, which could not be compared with the work of other authors. To circumvent the above condition, the Qd quantities are found by calculating separately for the temperature intervals of 750-850 and 950-1050 °C. The plots of ln strain rate versus ln (sinh (ar)),

Table 5 Material constants of steels calculated using peak stress values, obtained from hot compression tests Temperature, °C

A, s21

n

Qd, kJ/mol

750-850 950-1050 750-850 950-1050

6.9 9 1013 1.9 9 108 3.0 9 1014 2.9 9 105

4.48 4.48 3.88 3.88

319.4 242.9 328.9 172.8

S1 S1 S2 S2

0

ln (strain rate)

-1

-2

750 a 800 b 850 c 900 d 950 e 1000 f 1050 g Linear fit

1.2

c

g b

0.8

f e a

R a (0.99) b (0.99) c (0.98) d (0.99) e (0.99) f (0.98) g (0.99)

-3

-4

0.2

0.4

a

0.2 0.0

2

R a (0.98) b (0.99) c (0.93)

-0.2 -0.4

0.4

0.6

0.8

1.0

1.2

ln sinh(ασ)

(a)

8.8

9.0

9.2

1.0 0.8

0.8

ln sinh(ασ )

c

0.4 0.2 b

0.0

9.8

0.01s–1 a 0.1 s–1 b 1s–1 c Linear fit

1.0

0.6

9.6

10000/T (K )

1.2

a 0.01s–1 b 0.1 s–1 c 1s–1 Linear fit

9.4 -1

(b)

1.2

ln sinh(ασ )

b

-0.6

-5 -0.6 -0.4 -0.2 0.0

c

b

0.6 0.4 0.2

a

0.0

2

-0.4

c

0.6 2

-0.2

0.01s–1 a 0.1s–1 b 1s –1 c Linear fit

1.0

d

ln sinh(ασ)

Steel

R a (0.94) b (0.99) c (0.99)

2

R a (0.93) b (0.89) c (0.98)

-0.2

a

-0.4 8.1

(c)

8.2

8.3

8.4

8.5

8.6

8.7

8.8

8.9

9.0

-1

10000/T (K )

7.5

(d)

7.6

7.7

7.8

7.9

8.0

8.1

8.2

-1

10000/T (K )

Fig. 10 (a) Relationship between ln (strain rate) and ln (sinh (ar)); relationship between ln (sinh (ar)) and 10,000/T for the temperature interval: (b) 750-850 °C, (c) 850-950 °C, and (d) 950-1050 °C; a = 0.012 for S1 steel

and ln (sinh (ar)) versus 10,000/T are displayed in Fig. 10 and 11 for S1 and S2 steels, respectively. The hot working tests were done over a wide range of steel phases, starting with complete ferrite phase, then with mixed (a + c) phases and finally in the purely austenitic phase region. This could be the reason for the low R2 values obtained in some linear fittings. The activation energy of S1 steel is found to be higher than S2 steel for the temperature interval of 9501050 °C. The n values between 3.8 and 4.5 for sinh law are consistent with the dislocation creep behavior (Ref 38). The deformation mechanism is dislocation climb controlled at low strain rates

and dislocation glide controlled at moderate and high strain rates (Ref 39). In the case of solid solution alloys, the solute atoms might preferentially segregate around moving dislocations in the form of Cottrell or impurity atmospheres. This can create a situation where the rate of dislocation glide becomes slower than the rate of dislocation climb. Thus, creep is controlled by glide in the form of a viscous drag process in the intermediate and high stress situations (Ref 40-43). Since the resistance to dislocation movement increases when alloying elements are added to the alloy, this addition lead to an increase in Qd (Ref 44). Furthermore, the addition of silicon to a carbon steel is expected to result in higher flow stress and a higher

Journal of Materials Engineering and Performance

2.0

2

ln (strain rate)

-1

-2

R a (0.98) b (0.99) c (0.99) d (0.99) e (0.99) f (0.99) g (0.99)

c

g

d

e f

-3

750 800 850 900 950 1000 1050 Linear fit

b

-0.3

0.0

0.3

0.6

0.9

1.2

1.5

a b c d e f g

0.5 a

-0.5 8.8

9.0

9.2

1.0

2

R a (0.55) b (0.97) c (0.94)

ln sinh(ασ )

0.8 0.6

9.8

0.01s –1 0.1s –1 1s Linear fit

a b c

c 2

R a (0.56) b (0.88) c (0.99)

0.8

0.2 b

0.0

9.6 -1

1.0

c

0.4

9.4

10000/T (K ) 1.2

a b c

ln sinh(ασ )

–1

c

b

(b)

0.01s –1 0.1s –1 1s Linear fit

2

R a (0.99) b (0.99) c (0.96)

1.0

–1

1.2

a b c

0.0

1.8

ln sinh(ασ)

(a)

0.01s –1 0.1s –1 1s Linear fit

1.5

-4

-5 -0.6

–1

a

ln sinh(ασ )

0

0.6 b

0.4 0.2 0.0

a

-0.2 -0.2 a

-0.4 8.1

8.2

(c)

8.3

8.4

8.5

8.6

8.7

8.8

8.9

-0.4 7.5

9.0

10000/T (K-1 )

7.6

7.7

7.8

7.9

8.0

8.1

8.2

10000/T (K-1 )

(d)

Fig. 11 (a) Relationship between ln strain rate and ln (sinh (ar)); relationship between ln (sinh (ar)) and 10,000/T for the temperature interval: (b) 750-850 °C, (c) 850-950 °C, and (d) 950-1050 °C; a = 0.012 for S2 steel

Table 6 Z values for the steels at 1, 0.1, and 0.01 s21 and at 800 and 1050 °C Temperature, °C 800 800 800 1050 1050 1050

Strain rate, s21 1 0.1 0.01 1 0.1 0.01

S1 steel Z 7.1 7.1 7.1 7.2 7.2 7.2

9 9 9 9 9 9

1020 1019 1018 1015 1014 1013

activation energy as compared to carbon steel in high Z conditions (Ref 45, 46). Medina and Hernandez (Ref 34) have reported similar findings using torsion tests on low alloy steels at various temperatures and strain rates. Addition of silicon and manganese increases the activation energy of deformation of low alloy steels (Ref 34). The activation energy of S2 steel is found to be higher than S1 steel for the temperature interval 750-850 °C. This implies that the higher alloying in S2 steel causes an increase in flow stress. The corrected peak flow

Journal of Materials Engineering and Performance

Mean grain size, lm 5.5 5.8 7.6 11.5 13.1 13.4

± 0.5 ± 0.1 ± 0.1 ± 1.6 ± 1.9 ± 1.9

S2 steel Z 2.9 2.9 2.9 1.9 1.9 1.9

9 1021 9 1020 9 1019 9 1011 9 1010 9 109

Average grain size, lm 6.7 6.5 9.6 12.9 13 13

± ± ± ± ± ±

0.1 0.8 0.7 1.6 1 2.6

stresses of S1 steel are lower than those of S2 steel at deformation temperatures of 750, 800, 900, 950, and 1000 °C using a strain rate of 1 s1 and at 800, 850, 1000, and 1050 °C using a strain rate of 0.1 s1, and at 850 and 1000 °C using a strain rate of 0.01 s1 (Table 4 and Fig. 4b and d). High values of Z result in grain refinement of steels during hot compression (Ref 18, 47, 48). High values of Z also suggest higher flow stresses for S2 steel (Ref 49). The average grain size decreases from 12.9 ± 1.6 lm for S2 steel to

11.5 ± 1.6 lm for S1 steel, at deformation conditions of 1050 °C and 1 s1 (Table 6) because S1 steel exhibits higher Z at this temperature and strain rate.

4. Conclusions The microstructural evolution and ferrite grain refinement of two high phosphorous steels are examined under a number of strain rates and deformation temperatures using hot compression tests. The following are the main conclusions: 1. Average grain size of 5-7 lm can be obtained for both steels when worked at lower temperatures in the intercritical (a + c) region. 2. Higher alloying in S2 steel increases the activation energy, resulting in the increase in flow stress in the temperatures range of 750-850 °C. 3. The grain size decreases with an increase in strain rate from 0.01 to 0.1 s1 and with a decrease in deformation temperature. 4. The n values obtained for these steels indicate that dislocation creep is the deformation mechanism.

Acknowledgments Authors are grateful to Vaishnav Steel Private Limited, Muzaffarnagar, India, for providing the steel castings for research purposes.

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