Dynamic recrystallization and precipitation in high ... - Springer Link

International Journal of Minerals, Metallurgy and Materials Volume 21, Number 1, January 2014, Page 36 DOI: 10.1007/s12613-014-0862-4

Dynamic recrystallization and precipitation in high manganese austenitic stainless steel during hot compression Amir Momeni1), Shahab Kazemi2), Golam Ebrahimi3), and Alireza Maldar 3) 1) Materials Science and Engineering Department, Hamedan University of Technology, Hamedan 3733-1-65169, Iran 2) Department of Materials Engineering, Faculty of Engineering, Bu Ali Sina University, Hamedan 4161-65174, Iran 3) Department of Materials and Polymer Engineering, Hakim Sabzevari University, Sabzevar 9617976487, Iran (Received: 18 July 2013; revised: 3 September 2013; accepted: 9 September 2013)

Abstract: Dynamic recrystallization and precipitation in a high manganese austenitic stainless steel were investigated by hot compression tests over temperatures of 950-1150°C at strain rates of 0.001 s−1-1 s−1. All the flow curves within the studied deformation regimes were typical of dynamic recrystallization. A window was constructed to determine the value of apparent activation energy as a function of strain rate and deformation temperature. The kinetics of dynamic recrystallization was analyzed using the Avrami kinetics equation. A range of apparent activation energy for hot deformation from 303 kJ/mol to 477 kJ/mol is obtained at different deformation regimes. Microscopic characterization confirms that under a certain deformation condition (medium Zener-Hollomon parameter (Z) values), dynamic recrystallization appears at first, but large particles can not inhibit the recrystallization. At low or high Z values, dynamic recrystallization may occur before dynamic precipitation and proceeds faster. In both cases, secondary phase precipitation is observed along prior austenite grain boundaries. Stress relaxation tests at the same deformation temperatures also confirm the possibility of dynamic precipitation. Unexpectedly, the Avrami's exponent value increases with the increase of Z value. It is associated with the priority of dynamic recrystallization to dynamic precipitation at higher Z values. Keywords: austenitic stainless steel; hot deformation; dynamic recrystallization; precipitation; activation energy

1. Introduction The microstructure evolution of metallic materials during hot deformation has the practical importance due to its relationship with mechanical properties of final products. Several softening processes, such as dynamic recovery (DRV), dynamic recrystallization (DRX), static recrystallization (SRX), and metadynamic recrystallization (MDRX), can influence the microstructures of deformed materials [1-5]. In some industrial hot working processes, such as rapid rod rolling or hot forging, the strain value is large enough to trigger DRX in austenitic materials due to low stacking fault energy (SFE) [6]. The influences of DRX on microstructure, flow behavior, and instability have been the subject of many investigations in recent years [7-9]. The kinetics and extent of DRX strongly depend on two Corresponding author: Amir Momeni

variables, material characteristics and processing parameters. In some materials, secondary phase precipitation may be stimulated by the deformation and interaction with DRX. Because both precipitation and recrystallization consume the stored energy of deformation, the occurrence of dynamic precipitation or phase transformation preceding DRX may use up the deformation energy, retarding or inhibiting the occurrence of recrystallization. The interactions between strain-induced precipitation (SIP) and SRX as well as dynamic precipitation (DP) and DRX in microalloyed steels have been studied extensively [10-14]. The resistance of DP on DRX is not only from the consumption of deformation energy stored in materials but also from the pinning effect of high-angle boundaries (HAGBs) by fine secondary phase particles. The pinning effect of precipitation on HAGBs was investigated and modeled in recent years by Xu et al. [11]. Jonas [15] also indicated that

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A. Momeni et al., Dynamic recrystallization and precipitation in high manganese austenitic stainless steel during hot …

when the particles gradually coarsened during deformation and lost their pinning force, DRX might resume and proceed at a lower speed. Results obtained from the released literature confirm that an in-depth knowledge about the interaction between DRX and DP is crucial to the proper design of hot working processes for different alloys. The existence of Cr (as a strong carbide former element) and Mn, which often cooperate in the formation of carbides, makes AISI 200 series of stainless steels prone to DP potentially. This category of austenitic stainless steels is an alternative to AISI 300 series, due to the lower cost, higher strength, and moderate corrosion resistance [16]. Although dynamic precipitation in high Cr stainless steels has been observed in stainless steels [17], their capability to interact with DRX, when other elements, such as Nb, are absent, is still in doubt [18]. On the other hand, the formation of Cr carbides increases the tendency of pitting and general corrosion because of Cr depletion that often occurs around grain boundaries [19]. Therefore, the research on the formation of Cr-carbides during hot working helps us to avoid pitting and general corrosion or design an appropriate post-deformation heat treatment to provide a single phase austenitic structure for a better corrosion resistance. Hence, the aim of the present work was to analyze the deformation, recrystallization, and precipitation behaviors of AISI 202 stainless steel to determine the best processing condition and find a better working response of final products.

2. Experimental procedures The material used in this investigation was AISI 202 high-manganese austenitic stainless steel with the composition of 0.12% C, 18.1% Cr, 0.7% Si, 8.5% Mn, 0.03% P, 0.01% S, 5.1% Ni, and the balance of Fe by mass. Cylindrical compression samples of 10 mm in diameter and 15 mm in height were prepared from the as-received bar whose initial microstructure is shown in Fig. 1. Concentric grooves of 0.5 mm in depth were machined on the surfaces of compression samples to keep the lubricant (boron nitride) in the contacting surfaces. A 1-mm chamfer was machined at an angle of 45° on the specimen edges to avoid the fold-over of the specimen at the early stages of hot compression testing. A Zwick-Roell 250 testing machine, equipped with a fully computerized furnace, was used to perform hot compression tests. Before testing, all the specimens were reheated to 1200°C, held for 15 min, and then cooled down to the testing temperature. Continuous hot compression tests were carried out at temperatures of 950°C to 1100°C with the intervals of 50°C and at strain rates from 0.001 s−1 to 1 s−1.

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After a total true strain of 0.7, the deformed samples were immediately quenched in water, cut along the longitudinal axis, and then prepared with the standard metallographic techniques. The electrolytic etching of 60vol% HNO3 solution was used to reveal the outline of microstructure.

Fig. 1. Starting microstructure of AISI 202 stainless steel used in this investigation.

Stress relaxation tests were conducted after the same reheating regime mentioned above. The technique was proved to be reliable in revealing the start and finish time of precipitation, the strain-induced precipitation, and the interactions between recrystallization and precipitation [20]. The constant pre-strain for stress relaxation was adopted as 0.05 to avoid the static recrystallization when stress would be monitored. Therefore, the variation of stress with time reflected the dynamic recovery and likely the precipitation effect which was very analogous to the onset of hot compression tests.

3. Results and discussion 3.1. Flow curves and constitutive analysis True stress-strain curves obtained from the load-stroke data are shown in Fig. 2. The curves mostly exhibit a single peak that is often attributed to the occurrence of DRX during hot deformation. From the start of testing up to the peak, DRV and work hardening occur concomitantly. When the dislocation density reaches a critical value, new grains nucleate and lead to flow softening. The steady-state condition is attained when the material is fully recrystallized and a balance is established between DRX and work hardening. Flow curves in Fig. 2 clearly indicate that the peak point and the steady-state condition are shifted to higher strains when the temperature declines or the strain rate rises. These effects clearly show that the kinetics of DRX directly depends on the deformation condition. Thus, the general stress-strain

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Int. J. Miner. Metall. Mater., Vol. 21, No. 1, Jan. 2014

behavior suggests that DRX occurs in a wide temperature range from 950°C to 1100°C and at strain rate ranging from 0.001 s−1 to 1 s−1. It is observed that the effects of temperature and strain rate on flow stress are significant for all testing conditions. The typical influences of temperature and strain rate on flow stress are observed clearly. To correlate the peak flow stress to the deformation condition, a hyper-

bolic sine constitutive equation was given as below. n ⎛ Q ⎞ Z = ε&exp ⎜ (1) ⎟ = A ⎡⎣sinh(ασ p ) ⎤⎦ ⎝ RT ⎠ where Z is Zener-Hollomon parameter, ε& the strain rate, Q the apparent activation energy for hot deformation, R the gas constant that is equal to 8.314 J/(mol·K), T the temperature, σp the peak stress, and A, n, and α the material constants.

Fig. 2. Flow curves obtained under different deformation conditions: (a) 950°C; (b) 1000°C; (c) 1050°C; (d) 1100°C; (e) 1150°C.

Eq. (2) is obtained by taking logarithm for both sides of Eq. (1). lnε& +

Q = ln A + n ln ⎡⎣sinh(ασ p ) ⎤⎦ RT

(2)

By substituting the experimental values of strain rate and flow stress (at the peak points) into Eq. (2), the dependence of ln[sinh(ασp)] on ln ε& and 1/T is plotted as shown in Fig. 3. α, the stress multiplier, is an adjustable constant, which is


determined by try and error to adjust the curves in the most linear and parallel condition, and defined as 0.014 here. The strain rate sensitivity parameter (m) is often defined as ∂lnσ/∂ln ε& [21], but according to Eq. (2), as (1/n) is equal to ∂ln[sinh(ασp)]/∂ln ε& , it can be taken as an index of the sensitivity of flow stress to strain rate. It is approved that there is a direct relationship between the strain rate sensitivity parameter (1/n) and the hot workability of materials. Hence, according to the variation of strain rate sensitivity

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parameter with deformation temperature in Fig. 4, it is deduced that the hot workability degrades at ~1050°C. A better workability at 950°C and 1000°C is observed, which is contrary to the general expectation as getting a better workability at higher temperatures [22-23]. This hot workability wave trough needs to be further addressed and analyzed by upcoming flow analysis and microstructural characterization. Eq. (2) is also used to determine the apparent activation energy as the following equation.

Fig. 3. Variation of peak flow stress with strain rate (a) and reciprocal of temperature (b) in the frame of the hyperbolic sine constitutive equation (S stands for the slope of curve, and Savg the average of slope).

determines how easy DRX starts at a given deformation condition (temperature and strain rate). Therefore, the window of Q values as a function of strain rate and temperature that is prepared by the same method described by Eqs. (1) to (3) can be used to evaluate the feasibility of DRX in the studied alloy. Fig. 5 provides the window to determine the value of Q at any strain rate and deformation temperature. This window indicates that at the temperature range of 1000-1100°C and low strain rates of 0.001 s−1 and 0.01 s−1,

Fig. 4. Strain rate sensitivity (1/n) as a function of deformation temperature.

Q = nR

{

∂ ln ⎡⎣sinh(ασ p ) ⎤⎦ ∂ (1/T )

}

(3)

The apparent activation energy at the peak point reflects the required energy to trigger DRX and flow softening. Therefore, it particularly depends on the chemical composition, deformation condition, and microstructure during hot working. The average value of Q for the studied material can be determined as 388 kJ/mol by multiplying the average slope (Savg=1.11) in Fig. 3(b) by the average value of n (=1/Savg= 4.2) in Fig. 3(a) and R. According to Eq. (3), the value of Q

Fig. 5. Window of Q values (the required energy to reach the peak point of the flow curve) at the different strain rates and deformation temperatures.

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the value of Q is unexpectedly high and decreases as the strain rate rises. At temperatures beyond this critical range, e.g., 1150°C, and low strain rates, the tendency for DRX increases so that it can be triggered at low activation energies. The value of Q increases when a competing microstructural phenomenon, such as DP, consumes the energy that is pumped into the material by deformation. Dynamic precipitation may be delayed or completely prohibited, depending on the nature of precipitates. In microalloyed steels, it has been well-understood that very fine (actually nanosized) carbides of microalloying elements pin grain boundaries and effectively inhibit DRX during hot deformation [10-14]. As a rule of thumb, it is believed that precipitates with the av-

erage diameter about 1 μm and larger often accelerate the recrystallization through the particle stimulating recrystallization (PSN) mechanism, while the finer ones may delay or inhibit the recrystallization [6]. Fig. 6 reflects the variation of peak stress (σp) and strain (εp) as functions of Z parameter. The practical importance of these curves is to predict the starting point of flow softening at any deformation condition and to generalize the lab results to the industrial applications. It is manifest that both σp and εp increase with increasing Z (increasing strain rate and/or decreasing deformation temperature). Indeed, the results confirm that DRX commences at a higher stress and strain when the temperature declines or the strain rate rises.

Fig. 6. Variation of peak strain (a) and peak stress (b) with Zener-Hollomon parameter (Z).

3.2. Microstructural characterization The microstructure of the sample after reheating at 1200°C for 15 min is shown in Fig. 7. This observation confirms that at the start of the prescribed thermomechanical regimes, the sample is a single-phase austenitic alloy with the average grain size of 60 μm and free of the remarkable secondary phase particles. The micrographs of hot deformed samples at the typical strain rate of 0.001 s−1 are observed in Fig. 8. By comparing Figs. 7 and 8, it can be found out that the precipitates present at all deformation temperatures in Fig. 8 should be formed during straining. Due to the coexistence of Cr and C with high content, it is reasonable to assume that the observed particles are chromium carbides that are stimulated by the influence of deformation energy. The equilibrium phase diagram of the studied alloy in Fig. 9 confirms that the particles should be the complex carbides composed of Cr, Mn, Fe, and C atoms. These results imply that the possible interaction of DP and DRX should be considered in discussing the hot deformation behavior of the studied alloy. One range of the crucial temperature in Fig. 8 is 950-1100°C (as shown in Figs. 8(a)-(d)). In this tempera-

ture range, grain boundary precipitates start and develop, and the particles become more or less coarse. Whereas at 1150°C (Fig. 8(e)), coarse precipitates are observed in the grain interiors as well as grain boundaries. It is also worthy of attention that the grain size gradually increases from 950°C to 1100°C; but at 1150°C, the grains are remarkably refined. With regard to the DRX flow curves presented in Fig. 2, it is deduced that the vast precipitation observed at all temperatures (as shown in Fig. 8) has not effectively suppressed DRX. This is an important finding that needs to be carefully probed by studying the microstructural changes and the evolution of particles size and place. Current observations indicate that the interaction of DP and DRX strongly depends on deformation temperature. Figs. 8(a)-(d) suggest that at the temperatures of 950-1100°C, DP precedes DRX, and therefore, the particles appear at prior austenite grain boundaries. As the average size of carbide particles is more or less over 1 μm, they cannot pin the boundaries, and therefore, DRX proceeds. The segregation of Cr and C atoms to prior austenite grain boundaries before deformation is a crucial step before DP. The segregation strongly depends on temperature. At very high temperatures, e.g.,


1150°C, the tendency for DRX increases due to the higher mobility of grain boundaries. At this temperature, the solute atoms of Cr and C cannot diffuse as fast as the boundaries, and therefore, DP lags behind DRX. It has been well described previously that the segregation of solute atoms to grain boundaries and the corresponding solute drag effect degrade beyond a critical temperature [24]. In such a case, DRX precedes DP as shown at 1150°C in Fig. 8(e), the precipitation actually occurs at the boundaries of dynamically formed grains. In Fig. 8(e), it is also interesting to note that particles observed inside the grains actually form at prior austenite grain boundaries and are a bit coarser than the grain boundary particles. In summary, at 0.001 s−1 and 950-1100°C, DP precedes DRX, and the particles decorate prior austenite grain boundaries; at 0.001 s−1 and 1150°C, DRX precedes DP, and the particles decorate the boundaries

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of DRX grains. It is notable that as the average size of the particles is more or less about 1 μm, no PSN is observed in the micrographs.

Fig. 7. Microstructure of the studied material after reheating at 1200°C for 15 min.

Fig. 8. Micrographs of the hot deformed samples at the strain rate of 0.001 s−1: (a) 950°C; (b) 1000°C; (c) 1050°C; (d) 1100°C; (e) 1150°C.

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Fig. 9. Phase equilibrium of the studied alloy, representing the phases expected to be present within the hot working range.

These results about the interaction of DRX and DP are also in agreement with the calculated apparent activation energies in Fig. 5. The higher values of Q at the temperature range of 950-1100°C (419-447 kJ/mol) with respect to the average value of 388 kJ/mol can be simply attributed to the priority of DP to DRX. As mentioned, although such coarse particles cannot suppress DRX, they can delay it due to consuming a part of deformation energy. On the contrary, at 1150°C, DRX happens more easily before DP (lower Q value, actually 361-390 kJ/mol), and then, DP proceeds at a lower speed. It is therefore concluded that there is a good correlation between the microstructural observations and the

constructed Q-window in Fig. 5. Using this correlation, the provided discussion can be generalized to the other deformation conditions. According to Fig. 5, at 950-1100°C, the Q value decreases as the strain rate rises. This is due to the fact that the mobility of grain boundaries and the tendency for DRX increase with increasing strain rate. Typical micrographs in Fig. 10 confirm this discussion. At a higher strain rate (1 s−1), DRX is stimulated more than precipitation, and the highly mobile grain boundaries accommodate less alloying elements. Hence, after the start of DRX, when enough solute atoms are populated at prior grain boundary sites, DP comes into operation. Micrographs in Fig. 10 clearly indicate the population of fine carbide particles along prior austenite grain boundaries. In the previous work on superaustenitic stainless steel, the complex carbides of Cr, Fe, and Ni were found to be responsible for DP and the interaction with DRX [25]. More recently, Zeng et al. [18] showed that pure Cr carbide could not form dynamically in AISI 403Nb martensitic stainless steel containing about 0.22wt% Nb, and the delay in DRX was attributed to DP of some complex carbides containing Nb, Cr, Fe, and C. The review of these previous results reflected that in the present material, high Mn concentration provided a driving force for the formation of (Cr,Mn) complex carbides. This may be due to the mutual interaction of Mn and Cr for increasing their activity at prior austenite grain boundaries and thereby the initiating precipitation.

Fig. 10. Micrographs of the hot deformed samples at the strain rate of 1 s−1: (a) 950°C; (b) 1100°C.

3.3. Stress relaxation tests and kinetics of DRX The possibility and extent of DP can be analyzed by stress relaxation tests whose underlying mechanisms have been described in Ref. [20]. A pre-strain of ~5% was applied to the samples to induce precipitation. This pre-strain is taken well below the critical strain for SRX, and therefore, in the absence of SRX, any change of stress that is monitored can be attributed to the strain induced precipitation

(SIP). Fig. 11 provides the stress relaxation curves at different deformation conditions. When there is no possibility for the softening or strengthening phenomenon, the stress is expected to decrease gradually. In Fig. 11, the curves of 950°C and 1000°C indicate a delay in stress release in a period of relaxation time; however, at higher temperatures of 1050-1150°C, even secondary hardening is observed. The delay in stress release or secondary hardening is often due to the interaction of SIP with moving dislocations. It should be


noticed here that the kinetics of DP has been known to be an order of magnitude faster than SIP [26]. Therefore, stress relaxation tests are only useful to study the possibility of DP and measure the precipitation starting and finishing times roughly. The results of stress relaxation tests for the studied material confirm that DP happens during hot deformation within the studied deformation regime. It is also inferred that the potential of DP increases as the deformation temperature rises. However, as the tendency of recrystallization increases with increasing temperature, DRX becomes prior to DP at high temperature. It is worthy of note that the critical temperature beyond which DRX overcomes DP increases with increasing strain rate. This is clearly observed from low Q regions in Fig. 5 and micrographs in Figs. 8 and 10. Actually, the regions with low and high Q values correspond to the priority of DRX and DP, respectively.

Fig. 11. Stress relaxation curves after 5% pre-strain to understand the possibility of dynamic precipitation at different deformation temperatures.

Based on the theory of nucleation and growth, previous researchers used the Avrami's kinetics equation to describe the progression of DRX with strain as follows [25]:

X = 1 − exp[ − k (ε − ε c ) n ]

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(4)

where X denotes the fractional softening, k and n the material constants, and εc the critical strain for the initiation of DRX. The DRX volume fraction under the different deformation conditions can be determined by either quantitative metallographic measurement or flow curve analysis. Compared to quantitative metallographic measurement, flow curve analysis is a more simple method, as the former needs a large quantity of samples to be examined, and it can be difficult to ascertain new grains precisely [27]. Here, to determine the progress of dynamic recrystallization volume fraction (XDRX), the following expression is employed as σp −σ X = (5) σp −σs where σp and σs are the peak and steady-state stress, and σ the flow stress between σp and σs. To determine k and n, Eq. (4) can be written as the following equation: ⎡ ⎛ 1 ⎞⎤ (6) ln ⎢ ln ⎜ ⎟ ⎥ = ln k + nln(ε − ε c ) ⎣ ⎝ 1 − X ⎠⎦ The Avrami's exponent (n) that is determined by Eq. (6) can be considered as an index to the DRX completion rate. The exponent n has a direct influence on the rate of DRX and can be used to determine the rate at which the instant flow stress descends from σp to σs. Therefore, the value of n can be used to compare the recrystallization rate at different deformation conditions and distinguish the regimes controlled by DP or DRX. Fig. 12(a) represents the typical curves drawn based on Eq. (6) to determine the value of n. It is clearly observed that in the deformation regimes of 950°C and 0.001 s−1, 1050°C and 0.001 s−1, and 1100°C and 0.1 s−1 that all belong to high Q-value regions in Fig. 5, the values of n (slope of the curves) are lower as compared to n in the regime of 1000°C and 0.01 s−1. At the latter deformation

Fig. 12. Typical Avrami graphs of DRX kinetics to determine the exponent n as the slopes at different deformation conditions (a) and the variation of exponent n with the Zener-Hollomon (Z) parameter (b).

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regime, according to Fig. 5, DRX starts before DP and therefore proceeds faster. The same conclusion is inferred in Fig. 12(b), where the variation of n with the Z parameter is drawn. It is generally expected that the value of n decreases with increasing Z, as previously reported for single phase austenitic alloys, such as 410 stainless steel [28]. This is associated with the decreasing tendency for DRX at higher strain rates and lower deformation temperatures. Unlikely for the present alloy, at low Z values where DP is prior to DRX, the value of n is relatively lower and increases with increasing Z values. These results are in accordance with Fig. 5 and indicate that, at high Z values when there is a lower tendency for the segregation of solute atoms to prior austenite grain boundaries and therefore a lower tendency for precipitation, DRX proceeds at higher speeds.

4. Conclusions Dynamic recrystallization and precipitation in a high manganese austenitic stainless steel were investigated. The most important results are as follows. (1) All the flow curves within the studied ranges of strain rate and temperature are typical of dynamic recrystallization. (2) A range of Q-values from 303 kJ/mol to 477 kJ/mol is obtained at different deformation regimes. The lowest Q-values were obtained at very high temperatures and low strain rates, or vice versa. (3) At a certain deformation condition (medium Z values), dynamic precipitation precedes dynamic recrystallization. However, large precipitation particles are not able to inhibit recrystallization. At low or high Z values, dynamic recrystallization occurs before dynamic precipitation and proceeds faster. In both cases, the precipitation of secondary phase particles is observed along prior austenite grain boundaries. (4) Stress relaxation tests at the same deformation temperatures confirm the possibility of dynamic precipitation. (5) The kinetics of dynamic recrystallization increases unexpectedly with increasing Z parameter. It is associated with the priority of dynamic recrystallization to dynamic precipitation at higher Z values.

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Acknowledgements The authors are grateful for the financial support from Hamedan University of Technology (No. 16.91.294).

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