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Study on hot workability and optimization of process parameters of a modified 310 austenitic stainless steel using processing maps
Materials and Design 67 (2015) 165–172
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Study on hot workability and optimization of process parameters of a modified 310 austenitic stainless steel using processing maps Hongying Sun a,c, Yongduo Sun b, Ruiqian Zhang b, Man Wang a, Rui Tang b, Zhangjian Zhou a,⇑ a
School of Materials Science and Engineering, University of Science and Technology Beijing, Xueyuan Road 30, Haidian District, Beijing 100083, China Science and Technology on Reactor Fuel and Materials Laboratory, Nuclear Power Institute of China, P.O. Box 436, Chengdu, Sichuan 610041, China c School of Mechanical Engineering, Anyang Institute of Technology, West of Huanghe Road, Wenfeng District, Anyang, Henan Province 455002, China b
a r t i c l e
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Article history: Received 6 September 2014 Accepted 25 November 2014 Available online 3 December 2014 Keywords: Austenitic stainless steel Hot workability Processing map Microstructure Hardness
a b s t r a c t To investigate the optimized hot deformation parameters of a modified 310 austenitic stainless steel, the hot compression tests were performed using a Gleeble 3500 thermal simulator. The hot deformation behavior and hot workability characteristics were investigated in a temperature range of 800–1100 °C and a strain rate range of 0.1–10 s1. The hot processing maps of the tested steel were developed based on the dynamic material model (DMM), from which the safe deformation regions and instable deformation regions were determined. The corresponding microstructural and hardness evolutions during deformation were analyzed in detail. It was found that the deformation in the safe regions was beneficial to dynamic recovery (DRY) and dynamic recrystallization (DRX), while the deformation in unstable region would lead to flow instability, kink boundaries and grain growth. Near 950 °C, the energy dissipation rates were unusually lower, and the hardness of the deformed sample exhibited a significant increase, as a result of strain-induced precipitation. Coupled with the microstructure analysis and processing map technology, the workability map was schematically plotted and the optimal working conditions were determined. Such conditions were: temperatures in the range of 1075–1100 °C and strain rates in the range of 0.5–1.7 s1. These conditions are critical to attain an excellent homogeneous microstructure with fine grains after deformation for the modified 310 austenitic stainless steel. Ó 2014 Elsevier Ltd. All rights reserved.
1. Introduction Austenitic stainless steels have excellent oxidation resistance and superior high temperature creep properties as compared to ferrite stainless steels. For these reasons, austenitic stainless steels have been selected as one of important structural materials for applications in extreme environments, such as advanced nuclear industries [1] and power plants [2]. Austenitic stainless steels with well-developed microstructures and desired properties are usually formed by reasonable processing. However, several unfavorable defects may be caused inevitably by some improper forming processes [3,4]. Therefore, the selection of reasonable hot working parameters plays an important role in the manufacturing process. During the past few decades, many scholars have attempted to optimize the processing parameters of alloys through experiments and theoretical modeling [5,6]. Their studies have contributed to the understanding of the mechanisms of hot deformation. But it was still difficult for researchers to achieve their optimization aims ⇑ Corresponding author. Tel./fax: +86 10 62334951. E-mail address: [email protected] (Z. Zhou). http://dx.doi.org/10.1016/j.matdes.2014.11.041 0261-3069/Ó 2014 Elsevier Ltd. All rights reserved.
until the concept of the dynamic materials model (DMM) was first proposed by Prasad in 1984 [7]. Processing maps are usually constructed based on the DMM. Processing maps can be used to control microstructures, determine optimal deformation parameters [8] and explain different deformation mechanisms [9]. At present, processing maps based on the DMM have been widely used because of its characteristics of convenience, accuracy and quickness, such as for non-ferrous alloys [10] and several stainless steels [11]. Considerable attention has been focused on the constitutive flow behaviors and workability of 304L, AISI 304 and 316L using processing maps established by the corresponding contour maps. To obtain the desired microstructure and satisfied properties, different deformation mechanisms have been conducted in a temperature range of 600–1200 °C and a strain rate range of 0.001– 100 s1. AISI304L steel experienced dynamic recrystallization (DRX) at temperature of 1150 °C and strain rate of 0.1 s1 [12], while AISI 304 steel underwent DRX at 1100 °C and 0.1 s1 [13], and 316L stainless steel occurred DRX at 1250 °C and 0.05 s1. In addition, the dynamic recovery (DRY) region of 316L steel was at 900 °C and 0.001 s1 [14]. Consequently, Venugopal research group
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carried out industrial validations of processing maps by press forging, hammer forging, rolling and extrusion of AISI 304, 316L and 304L austenitic stainless steels [15,16]. The results showed that the hot workability of austenitic stainless steels highly depended on their chemical compositions and initial microstructure. Besides the DRX and DRY domains, processing maps also exhibit several unavoidable instability regions, such as flow localization, dynamic strain aging (DSA), adiabatic shear bands, kink bands and intense deformation bands. These regions are related to the low efficiency of energy dissipation at high strain rates [17], and provide the guidance for avoidance of the deformation defects. Tan’s work [18] on the processing maps and hot workability of super 304H demonstrated that the large energy dissipation efficiency contributed to the occurrence of DRX, and the optimal deformation parameters of super 304H were at 1100 °C and the strain rate was more than 0.5 s1. These researches validate that processing maps are significant and necessary to predict the optimal deformation parameters and guide the subsequent plastic forming in industry. Among all of the existing austenitic stainless steels, AISI 310 steel is an advanced austenitic stainless steel with improved oxidation and corrosion resistance properties. At present, although considerable effort has been expended in the research of constitutive flow behavior and workability of common austenitic stainless steels, little attention has been focused on the effects of precipitation process during hot deformation. Moreover, few studies have investigated the processing maps and the workability characteristics of advanced 310 austenitic steel. In particular, these available studies lack an examination of the effects about dynamic precipitation on the processing maps and workability. In our previous work, hot behaviors and microstructures of a modified 310 austenitic stainless steel have been researched in order to determine the optimum deformation parameters [19]. However, it is difficult to obtain required information due to the complicated influences of second phases and precipitates that exist prior to or during the deformation. In the present study, we establish the processing maps of the tested steel based on the DMM, and further describe the workability maps by microstructural analysis. Then we determine the reasonable working perimeters of the modified 310 steel based on the microstructural analysis and workability maps. 2. Experimental procedures The investigated material was a modified 310 austenitic stainless steel which was fabricated by vacuum induction melting. The chemical composition of this tested steel was designed as: 25Cr, 20Ni, 0.1C, 0.65Si, 0.2Ti, 0.2Zr, 0.15W, 0.15V, and balance Fe (wt.%). Cylindrical hot compression specimens with a size of U 8 mm 15 mm were cut from the forged plate. The initial microstructure of the specimen showed a single-phase of austenite with a small quantity of twins. All high-temperature compression tests were conducted by using a Gleeble-3500 thermal mechanical simulator. To minimize the friction during deformation, a layer of high-temperature lubricant composed of MoS2 was daubed on the sample surfaces. Before testing, specimens were heated to 1200 °C at the rate of 10 °C s1 and held for 10 min. The specimens were then cooled to different test temperatures at the rate of 5 °C s1, and these specimens were held for 5 min. The deformation temperatures were 800, 900, 950, 1000 and 1100 °C, and their strain rates were 0.1, 1 and 10 s1, respectively. Most of the samples deformed at a constant strain rate to the true strain of 0.68, while two additional samples were respectively compressed to the strains of 0.34 and 1.2 at 1100 °C and 1 s1. After deformation, the samples were instantly quenched into water.
The samples were sectioned from the center for the microstructural analysis. The sectioned samples were then progressively ground to 2000 grids with SiC sand papers, and then they underwent electrolytic polishing. And specimens were etched for 10– 30 s using 0.5g FeCl3 + 6 mL HCl + 10 mL H2O as a solution to reveal the microstructures. Microstructural observations were performed using an optical microscope (OM, ZEISS Observer. A1m) and a scanning electron microscopy (SEM, FEI Quata 200) equipped with electron back-scatter diffraction (EBSD) detector and an energy dispersive spectrometer (EDS). Rockwell hardness and Vickers hardness were measured by different sclerometers (TIME TH320 and Leica VMHT 30 M, respectively). Every sample was measured more than three times, and the result was their average value. 3. Results and discussion 3.1. Stress strain curves Fig. 1 shows the true stress-strain curves at various strain rates and temperatures. Almost all of the flow curves reveal peak flow stresses and flow softening. But when the strain reaches 1.2 at the strain rate of 1 s1 and temperature of 1100 °C, the curve shows a steady state. In the initial stage of compressive deformation, the hardening plays a significant role and the stress greatly increases with the strain. When the strain goes critical value, DRX is activated. The stress continues to increase until it reaches a peak value, at the same time the stored energy come to the highest value. It is known that the driving force for DRX is the stored energy, thus DRX become more dominant with the increase of the stored energy. On the other hand, the stored energy is consumed as DRX occurred, resulting in the activation of flow softening, and the stress decreases abruptly. Along with the consumption of stored energy, the rate of DRX gradually decreases until full DRX is achieved (see Fig. 1(d), e = 1.2). This dynamic equilibrium between the softening and hardening in the matrix is very difficult to be obtained, because DRX occurs in this low stacking fault energy steel by an important reversion mechanism [20,21]. Thus, the characteristics of recrystallization (steady state) in curves are not observed till the strain reaches 1.2. The shapes of the stressstrain curves, work hardening, flow softening, steady state and oscillation behaviors are related to the specific hot working process and inherent deformation mechanisms [22]. Understanding of the mechanisms of strain hardening and softening mechanisms during hot working is important for the investigation of the hot working process. Therefore, further additional studies about the deformation mechanisms were performed, which will be discussed in the following sections. 3.2. Hot processing maps The constitutive behaviors of materials are usually studied through the investigation of processing maps based on DMM, which has been described in detail by Prasad and Sasidhara [23]. The energy dissipation during hot deformation consists of two portions: one is used for plastic deformation, and the other is consumed by structural changes. The ratio (g) of the energy dissipation that is consumed by structural changes to the total energy is plotted with a contour line as a function of strain rate and temperature. Thus, an energy dissipation map is developed. The efficient energy dissipation ratio g during hot working is [24]:
g ¼ 2m=ðm þ 1Þ
ð1Þ
In Eq. (1), m is the strain rate sensitivity, which can be denoted as [25]:
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Fig. 1. True stress–true strain curves of tested steel under different deformation conditions. (a) e_ ¼ 0:1 s–1; (b) e_ ¼ 1 s–1 [19]; (c) e_ ¼ 10 s–1; (d) T = 1100 °C, e_ ¼ 1 s–1.
m¼
@ log r e_ ; T @ log e_
ð2Þ
At a certain temperature, m is a function of the stress and strain rate. Prasad and Seshacharyulu and Seshacharyulu [26] have validated that flow instability may occur during hot working when the following is satisfied:
n ¼ @ lnðm=ðm þ 1ÞÞ=@ ln e_ þ m < 0
cient energy dissipation ratio (g) at the same deformation conditions gradually increases with strain, mainly because the deformation stored energy increases during this process. It is mentionable that the efficient energy dissipation experiences unusual changes when temperatures near 950 °C as compared with other temperature zones. The causes of this phenomenon will be discussed in Section 3.4.
ð3Þ 3.3. Corresponding microstructural evolutions
If n is negative, flow instability will occur. The instability parameter varies with temperature and strain rate. Energy dissipation ratio (g) satisfies Eq. (3), temperature and strain rate in the contour map constitutes an instability map. A processing map will be obtained by superimposing the instability map on the energy dissipation map. In short, the processing map consists of safe domains and unsafe domains, which present specific microstructural mechanisms. Based on the stress strain curves and the methods described above, two processing maps for the modified 310 steel were obtained by Mat-lab software. The results are shown in Fig. 2. The numbers in the contour lines represent energy dissipation percentage (g %) at their responding deformation conditions. The gray regions denote rheological instability domains. As shown in Fig. 2, the processing maps at various strains exhibit similar broad features, but they are different in the details. There are one valley area of energy dissipation ((1)) and two energy dissipation peak areas ((2) and (3)) in Fig. 2(a) (marked with sequence number). Contrastingly, two energy dissipation valley areas ((4) and (5)) and an energy dissipation peak zone ((6)) are visible in Fig. 2(b). Also, two instability ranges exist at the strain of 0.128, and three instability ranges are exhibited when strain increases to 0.342. It is important to note that the size of instability areas slightly increases with the increasing of strain, which indicates a relevant forming ability of unstable microstructures. Additionally, the effi-
Different regions in the maps indicate specific microstructural evolutions after the specimen were treated by corresponding processing parameters. A narrow unstable region is observed at a lower strain rate range of 0.1–0.3 s1 when the temperature is 800 °C. The corresponding microstructure is shown in Fig. 3(a) and denotes a characteristic of flow instability, which demonstrates that hot working under such conditions should be avoided. It is encouraging that two favorable regions are observed in Fig. 2(b). The first one is in the temperature range of 820–900 °C and a strain rate of 0.1 s1. The peak efficiency is 48%, and it exists at the temperature of 850 °C and the strain rate of 0.1 s1. The microstructure in Fig. 3(b) exhibits a favorable characteristic of DRY. The other favorable region is in the temperature range of 1000–1100 °C and a strain rate range is confined to be less than 0.2 s1. In this region, the peak efficiency is 59% at 1075 °C and the strain rate of 0.1 s1. Accordingly, Fig. 3(c) presents the microstructural features that deformed in this range, in which DRX partially occurred. The microstructural feature of Fig. 3(c) is that some grains are DRX grains (black arrows-necklace structure), and others are deformed grains (other black arrows-pancaked grains). At the lowest strain rate and the highest temperature, the tested steel has almost completely undergone DRX, as seen in Fig. 3(d). Actually, some DRX grains have been proven to grow because of the peak efficiency. This result fits well with current literature [12],
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Fig. 2. Hot processing maps of tested steel at different true strains. (a) e = 0.128; (b) e = 0.342.
Fig. 3. Microstructures of the deformed specimens at different temperatures (arrows in (c) illustrate recrystallized grains and deformed grains). (a) 800 °C, 0.1 s–1; (b) 900 °C, 0.1 s–1; (c) 1000 °C, 0.1 s–1; (d) 1100 °C, 0.1 s–1.
in which DRX occured at the peak efficiency, and was followed by grain coarsening. Hot cracks may be caused in these coarse grains when the material is welded, so it is considered to be unstable. At higher strain rates, every map has two large unstable regions. One is in the lower temperature range, and the other is in the higher temperature range. As strain increases, the unstable areas become more extended and the DRY ranges will be much narrower. Flow instability microstructures such as shear bands tend to be formed at lower temperatures and higher strain rates. In the present investigation, the imposed strain is lower, so no unwanted defects are observed except for common flow instability. Though DRX has almost completely occurred at higher temperature (Fig. 4(b)), the microstructures display a distinct feature of kinked boundaries under this condition (see Fig. 4). The relevant flow stress fluctuates, as demonstrated by its ‘‘wavy’’ behavior,
seen in Fig. 1(c). These kink boundaries are particularly prone to intercrystalline cracking at elevated temperature, and are considered to be unstable. In addition, samples deformed at this elevated temperature range may introduce d-ferrite in uniform austenitic matrix. d-ferrite has deleterious effects on properties, which have been verified in AISI 310 austenitic stainless steel [27]. This result suggests that deformation in this region of the steel should be avoided. Based on a detailed analysis, it can be deduced that the domain at a medium-strain rate and higher temperature is beneficial to the microstructures. Fig. 5(a) [19] shows the microstructure corresponding to the point marked A, which is denoted by a black arrow in Fig. 2(b). Fig. 5(b) and (c) [19] show the microstructures of the samples that deformed at the same temperature and strain rate, but to the larger strains of 0.68 and 1.2, respectively. Fig. 5(b)
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Fig. 4. Microstructures of specimens deformed at different strain rates at 1000 °C and 1100 °C. (a) 1000 °C, e_ ¼ 10 s1; (b) 1100 °C, e_ ¼ 10 s1.
Fig. 5. EBSD images of the samples deformed at 1100 °C and 1 s1 to different strains. (a) e = 0.34 [19]; (b) e = 0.68; (c) e = 1.2 [19].
and (c) show the regular homogeneous austenitic microstructures, and they are different in the degrees of DRX. As seen in Fig. 5(c), DRX varies with the increase of strain, until DRX has been fully occurred. A finer and more uniform microstructure is obtained at the highest strain of 1.2, which supports the claim that this region provides an appropriate forming regime for obtaining homogeneous microstructures. A paradox is discovered between the stress-strain curve of Fig. 1(d) and the microstructure of Fig. 5(b). The former denotes DRY characteristics when the strain is equal to 0.68, but the latter exhibits the trait of DRX. This phenomenon supports the earlier deduction that DRX occurs in the austenitic stainless steel is a
dominant DRY mechanism [20]. Therefore, it is unreasonable to determine whether or not DRX occurred only by the analysis of stress-strain curves. 3.4. Hardness of the deformed sample As shown in Fig. 6, both Vickers hardness and Rockwell hardness of the tested steel gradually reduce with the temperature increase. At the strain rate of 1 s1, the highest hardness is 101HRB, while the lowest one is 76.6 HRB. The reduction percentage is about 24.4%. Similarly, at a strain rate of 0.1 s1, the difference in micro-hardness between the maximum and minimum values is 64 Hv. And the
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Fig. 6. The hardness of tested steel as a function of deformation temperature.
reduction ratio is 23.9%. It is interesting that the degrees of softening measured in the microstructures are roughly equal at these two strain rates. In Fig. 6, both hardness-temperature curves display significant increases near 950 °C. The deviation from the usual value will be larger when the strain rate increased. In order to understand the reason for these changes in hardness, we observed the microstructures of samples deformed at 950 °C. The microstructures show that the increase in hardness is related to the strain induced precipitation (SIP). Momeni et al. [28] studied the competition behavior between SIP and DRX for Fe–28Cr–33Ni–3.5Mo–0.6Mn steel and found that (Cr, Fe, Mo)23C6 precipitates nucleated at the certain temperature and lower strain rate [28]. The precipitates can be observed in the microstructures, in which weak gray GBs and black GBs are etched out, as designated with white and black
arrows in Fig. 7(a), respectively. Detailed examination of the black GBs, as shown in Fig. 7(b), shows that some fine pearl liked precipitates and a spot of white precipitates (indicated with white and black arrows in Fig. 7(b), respectively) are distributed on a high angle GB. Their chemical composition is analyzed by EDS, and the results indicate that they are secondary precipitates of (Cr, Fe, Mo)23C6 and Mo-rich intermetallic compounds of r phase (see Fig. 7(c) and (d)). The compositions in these two kinds of precipitates are very different. There are 5.05C, 73.71Cr, 16.38Fe, 3.87Mo (wt.%) in (Cr, Fe, Mo)23C6 carbides, and 0.28C, 36.31Cr, 44.60Fe, 6.98Mo (wt.%) in sigma precipitates. The precipitation and DRX are diffusion-controlled processes. SIP and DRX have certain inherent connections that may interact and restrict each other. If the temperature is too low, the diffusion rate of solute atoms is low, and precipitation is hard to occur, such as 800 °C. At lower temperatures, recrystallization process is difficult to occur, or may be accompanied by the formation of precipitates, because the solute atoms are liable to cluster at low deformed temperatures and strains [28]. At higher temperature, the diffusion rate of solute atoms is faster. And the pinning of high angle boundaries by SIP is diminished, so DRX is easy to occur. At the highest strain rate, i.e., 10 s1, the time for the occurrence of SIP is too short besides the higher driving force for DRX. Therefore, the DRX is dominated, which resulted in consuming the stored energy, and SIP is restricted. Similarly, during the precipitation, the stored energy in the system is consumed, and the energy dissipation becomes lower. Therefore, the precipitation is related to not only the temperature but also the strain rate. The interaction and restriction between SIP and DRX explain why the energy dissipation percentages near 950 °C are much lower. In this region, the plastic deformation is dominated. And the lager plastic deformation
Fig. 7. OM and SEM observations of precipitates at HAGBs and its EDS spectrum. (a) OM image of sample deformed 950 °C and 0.1 s1; (b) SEM observations of precipitates at HAGBs; (c) EDS spectrum of (Cr, Fe, Mo)23C6 precipitates; (d) EDS spectrum of r phase.
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(4) By the analysis of various regions in the processing maps and the observation of the corresponding microstructures, the workability map of the steel was schematically plotted, which more directly revealed the optimum hot working and instability ranges. The workability map indicated that the optimal deformation parameters of the tested steel were in the temperature range of 1075–1100 °C and the intermediate strain rate range of 0.5–1.7 s1.
Acknowledgements
Fig. 8. Hot workability map of the modified 310 steel under investigated conditions.
increases strain. The strain stimulates the precipitation. As a result, the DRX process is inhibited because of the lowered energy. Based on the analysis, it can be concluded that the precipitation occurred near 950 °C and lower strain rates. 3.5. Hot workability According to the above analysis, the hot workability (processing-microstructural mechanism) map was schematically plotted, which is shown in Fig. 8. For the modified 310 steel, it is obvious that the feasible hot working regime is in the DRX range. Therefore, the optimum hot working parameters of the steel are in the elevated temperature range of 1075–1100 °C and a strain rate range of 0.5–1.7 s1, which will be beneficial for optimizing the microstructures and mechanical properties of the tested steel. It should be noted that the range of appropriate working regimes are rather limited due to the evolution of the material (many coarse carbonization particles on triple joint GBs (Fig. 5, black areas)) and the effect of precipitation during the hot deformation. 4. Conclusions Based on the investigation of the hot processing maps and hot workability of a modified 310 austenitic stainless steel, and the temperature was in a range of 800–1100 °C and strain rate was in a range of 0.1–10 s1, some conclusions were drawn as follows: (1) The energy dissipation rate g and the hot processing maps of the modified 310 steel were obtained based on DMM. The instable deformation regions which should be avoided existed at a high strain rate, and in both low and high temperature ranges. Although the processing maps established at the different strains showed similar features, strain had a significant influence on the processing maps. (2) The deformation in the safe region was beneficial to DRY and DRX, while the deformation in unstable region would lead to flow instability, kink boundaries and grain growth. (3) The hardness of the deformed samples gradually reduced as the temperature increased except that the temperature near 950 °C, at which the hardness evidently increased due to the precipitation hardening. EDS spectroscopy analysis showed that the precipitates were secondary carbonizations of (Cr, Fe, Mo)23C6 and the intermetallic compound of r phase. In addition, because the stored energy was consumed by the precipitation, the SIP resulted in the decrease of energy dissipation rate (g) in the processing maps, and DRX was inhibited.
The authors are grateful for the financial support of Foundation of Science and Technology on Reactor Fuel and Materials Laboratory and the assistance of Xiao Li in discussing the EBSD images in this paper. The authors appreciate Miss. April Novak from University of Illinois for her contribution to the English revision.
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Materials and Design journal homepage: www.elsevier.com/locate/matdes
Study on hot workability and optimization of process parameters of a modified 310 austenitic stainless steel using processing maps Hongying Sun a,c, Yongduo Sun b, Ruiqian Zhang b, Man Wang a, Rui Tang b, Zhangjian Zhou a,⇑ a
School of Materials Science and Engineering, University of Science and Technology Beijing, Xueyuan Road 30, Haidian District, Beijing 100083, China Science and Technology on Reactor Fuel and Materials Laboratory, Nuclear Power Institute of China, P.O. Box 436, Chengdu, Sichuan 610041, China c School of Mechanical Engineering, Anyang Institute of Technology, West of Huanghe Road, Wenfeng District, Anyang, Henan Province 455002, China b
a r t i c l e
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Article history: Received 6 September 2014 Accepted 25 November 2014 Available online 3 December 2014 Keywords: Austenitic stainless steel Hot workability Processing map Microstructure Hardness
a b s t r a c t To investigate the optimized hot deformation parameters of a modified 310 austenitic stainless steel, the hot compression tests were performed using a Gleeble 3500 thermal simulator. The hot deformation behavior and hot workability characteristics were investigated in a temperature range of 800–1100 °C and a strain rate range of 0.1–10 s1. The hot processing maps of the tested steel were developed based on the dynamic material model (DMM), from which the safe deformation regions and instable deformation regions were determined. The corresponding microstructural and hardness evolutions during deformation were analyzed in detail. It was found that the deformation in the safe regions was beneficial to dynamic recovery (DRY) and dynamic recrystallization (DRX), while the deformation in unstable region would lead to flow instability, kink boundaries and grain growth. Near 950 °C, the energy dissipation rates were unusually lower, and the hardness of the deformed sample exhibited a significant increase, as a result of strain-induced precipitation. Coupled with the microstructure analysis and processing map technology, the workability map was schematically plotted and the optimal working conditions were determined. Such conditions were: temperatures in the range of 1075–1100 °C and strain rates in the range of 0.5–1.7 s1. These conditions are critical to attain an excellent homogeneous microstructure with fine grains after deformation for the modified 310 austenitic stainless steel. Ó 2014 Elsevier Ltd. All rights reserved.
1. Introduction Austenitic stainless steels have excellent oxidation resistance and superior high temperature creep properties as compared to ferrite stainless steels. For these reasons, austenitic stainless steels have been selected as one of important structural materials for applications in extreme environments, such as advanced nuclear industries [1] and power plants [2]. Austenitic stainless steels with well-developed microstructures and desired properties are usually formed by reasonable processing. However, several unfavorable defects may be caused inevitably by some improper forming processes [3,4]. Therefore, the selection of reasonable hot working parameters plays an important role in the manufacturing process. During the past few decades, many scholars have attempted to optimize the processing parameters of alloys through experiments and theoretical modeling [5,6]. Their studies have contributed to the understanding of the mechanisms of hot deformation. But it was still difficult for researchers to achieve their optimization aims ⇑ Corresponding author. Tel./fax: +86 10 62334951. E-mail address: [email protected] (Z. Zhou). http://dx.doi.org/10.1016/j.matdes.2014.11.041 0261-3069/Ó 2014 Elsevier Ltd. All rights reserved.
until the concept of the dynamic materials model (DMM) was first proposed by Prasad in 1984 [7]. Processing maps are usually constructed based on the DMM. Processing maps can be used to control microstructures, determine optimal deformation parameters [8] and explain different deformation mechanisms [9]. At present, processing maps based on the DMM have been widely used because of its characteristics of convenience, accuracy and quickness, such as for non-ferrous alloys [10] and several stainless steels [11]. Considerable attention has been focused on the constitutive flow behaviors and workability of 304L, AISI 304 and 316L using processing maps established by the corresponding contour maps. To obtain the desired microstructure and satisfied properties, different deformation mechanisms have been conducted in a temperature range of 600–1200 °C and a strain rate range of 0.001– 100 s1. AISI304L steel experienced dynamic recrystallization (DRX) at temperature of 1150 °C and strain rate of 0.1 s1 [12], while AISI 304 steel underwent DRX at 1100 °C and 0.1 s1 [13], and 316L stainless steel occurred DRX at 1250 °C and 0.05 s1. In addition, the dynamic recovery (DRY) region of 316L steel was at 900 °C and 0.001 s1 [14]. Consequently, Venugopal research group
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carried out industrial validations of processing maps by press forging, hammer forging, rolling and extrusion of AISI 304, 316L and 304L austenitic stainless steels [15,16]. The results showed that the hot workability of austenitic stainless steels highly depended on their chemical compositions and initial microstructure. Besides the DRX and DRY domains, processing maps also exhibit several unavoidable instability regions, such as flow localization, dynamic strain aging (DSA), adiabatic shear bands, kink bands and intense deformation bands. These regions are related to the low efficiency of energy dissipation at high strain rates [17], and provide the guidance for avoidance of the deformation defects. Tan’s work [18] on the processing maps and hot workability of super 304H demonstrated that the large energy dissipation efficiency contributed to the occurrence of DRX, and the optimal deformation parameters of super 304H were at 1100 °C and the strain rate was more than 0.5 s1. These researches validate that processing maps are significant and necessary to predict the optimal deformation parameters and guide the subsequent plastic forming in industry. Among all of the existing austenitic stainless steels, AISI 310 steel is an advanced austenitic stainless steel with improved oxidation and corrosion resistance properties. At present, although considerable effort has been expended in the research of constitutive flow behavior and workability of common austenitic stainless steels, little attention has been focused on the effects of precipitation process during hot deformation. Moreover, few studies have investigated the processing maps and the workability characteristics of advanced 310 austenitic steel. In particular, these available studies lack an examination of the effects about dynamic precipitation on the processing maps and workability. In our previous work, hot behaviors and microstructures of a modified 310 austenitic stainless steel have been researched in order to determine the optimum deformation parameters [19]. However, it is difficult to obtain required information due to the complicated influences of second phases and precipitates that exist prior to or during the deformation. In the present study, we establish the processing maps of the tested steel based on the DMM, and further describe the workability maps by microstructural analysis. Then we determine the reasonable working perimeters of the modified 310 steel based on the microstructural analysis and workability maps. 2. Experimental procedures The investigated material was a modified 310 austenitic stainless steel which was fabricated by vacuum induction melting. The chemical composition of this tested steel was designed as: 25Cr, 20Ni, 0.1C, 0.65Si, 0.2Ti, 0.2Zr, 0.15W, 0.15V, and balance Fe (wt.%). Cylindrical hot compression specimens with a size of U 8 mm 15 mm were cut from the forged plate. The initial microstructure of the specimen showed a single-phase of austenite with a small quantity of twins. All high-temperature compression tests were conducted by using a Gleeble-3500 thermal mechanical simulator. To minimize the friction during deformation, a layer of high-temperature lubricant composed of MoS2 was daubed on the sample surfaces. Before testing, specimens were heated to 1200 °C at the rate of 10 °C s1 and held for 10 min. The specimens were then cooled to different test temperatures at the rate of 5 °C s1, and these specimens were held for 5 min. The deformation temperatures were 800, 900, 950, 1000 and 1100 °C, and their strain rates were 0.1, 1 and 10 s1, respectively. Most of the samples deformed at a constant strain rate to the true strain of 0.68, while two additional samples were respectively compressed to the strains of 0.34 and 1.2 at 1100 °C and 1 s1. After deformation, the samples were instantly quenched into water.
The samples were sectioned from the center for the microstructural analysis. The sectioned samples were then progressively ground to 2000 grids with SiC sand papers, and then they underwent electrolytic polishing. And specimens were etched for 10– 30 s using 0.5g FeCl3 + 6 mL HCl + 10 mL H2O as a solution to reveal the microstructures. Microstructural observations were performed using an optical microscope (OM, ZEISS Observer. A1m) and a scanning electron microscopy (SEM, FEI Quata 200) equipped with electron back-scatter diffraction (EBSD) detector and an energy dispersive spectrometer (EDS). Rockwell hardness and Vickers hardness were measured by different sclerometers (TIME TH320 and Leica VMHT 30 M, respectively). Every sample was measured more than three times, and the result was their average value. 3. Results and discussion 3.1. Stress strain curves Fig. 1 shows the true stress-strain curves at various strain rates and temperatures. Almost all of the flow curves reveal peak flow stresses and flow softening. But when the strain reaches 1.2 at the strain rate of 1 s1 and temperature of 1100 °C, the curve shows a steady state. In the initial stage of compressive deformation, the hardening plays a significant role and the stress greatly increases with the strain. When the strain goes critical value, DRX is activated. The stress continues to increase until it reaches a peak value, at the same time the stored energy come to the highest value. It is known that the driving force for DRX is the stored energy, thus DRX become more dominant with the increase of the stored energy. On the other hand, the stored energy is consumed as DRX occurred, resulting in the activation of flow softening, and the stress decreases abruptly. Along with the consumption of stored energy, the rate of DRX gradually decreases until full DRX is achieved (see Fig. 1(d), e = 1.2). This dynamic equilibrium between the softening and hardening in the matrix is very difficult to be obtained, because DRX occurs in this low stacking fault energy steel by an important reversion mechanism [20,21]. Thus, the characteristics of recrystallization (steady state) in curves are not observed till the strain reaches 1.2. The shapes of the stressstrain curves, work hardening, flow softening, steady state and oscillation behaviors are related to the specific hot working process and inherent deformation mechanisms [22]. Understanding of the mechanisms of strain hardening and softening mechanisms during hot working is important for the investigation of the hot working process. Therefore, further additional studies about the deformation mechanisms were performed, which will be discussed in the following sections. 3.2. Hot processing maps The constitutive behaviors of materials are usually studied through the investigation of processing maps based on DMM, which has been described in detail by Prasad and Sasidhara [23]. The energy dissipation during hot deformation consists of two portions: one is used for plastic deformation, and the other is consumed by structural changes. The ratio (g) of the energy dissipation that is consumed by structural changes to the total energy is plotted with a contour line as a function of strain rate and temperature. Thus, an energy dissipation map is developed. The efficient energy dissipation ratio g during hot working is [24]:
g ¼ 2m=ðm þ 1Þ
ð1Þ
In Eq. (1), m is the strain rate sensitivity, which can be denoted as [25]:
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Fig. 1. True stress–true strain curves of tested steel under different deformation conditions. (a) e_ ¼ 0:1 s–1; (b) e_ ¼ 1 s–1 [19]; (c) e_ ¼ 10 s–1; (d) T = 1100 °C, e_ ¼ 1 s–1.
m¼
@ log r e_ ; T @ log e_
ð2Þ
At a certain temperature, m is a function of the stress and strain rate. Prasad and Seshacharyulu and Seshacharyulu [26] have validated that flow instability may occur during hot working when the following is satisfied:
n ¼ @ lnðm=ðm þ 1ÞÞ=@ ln e_ þ m < 0
cient energy dissipation ratio (g) at the same deformation conditions gradually increases with strain, mainly because the deformation stored energy increases during this process. It is mentionable that the efficient energy dissipation experiences unusual changes when temperatures near 950 °C as compared with other temperature zones. The causes of this phenomenon will be discussed in Section 3.4.
ð3Þ 3.3. Corresponding microstructural evolutions
If n is negative, flow instability will occur. The instability parameter varies with temperature and strain rate. Energy dissipation ratio (g) satisfies Eq. (3), temperature and strain rate in the contour map constitutes an instability map. A processing map will be obtained by superimposing the instability map on the energy dissipation map. In short, the processing map consists of safe domains and unsafe domains, which present specific microstructural mechanisms. Based on the stress strain curves and the methods described above, two processing maps for the modified 310 steel were obtained by Mat-lab software. The results are shown in Fig. 2. The numbers in the contour lines represent energy dissipation percentage (g %) at their responding deformation conditions. The gray regions denote rheological instability domains. As shown in Fig. 2, the processing maps at various strains exhibit similar broad features, but they are different in the details. There are one valley area of energy dissipation ((1)) and two energy dissipation peak areas ((2) and (3)) in Fig. 2(a) (marked with sequence number). Contrastingly, two energy dissipation valley areas ((4) and (5)) and an energy dissipation peak zone ((6)) are visible in Fig. 2(b). Also, two instability ranges exist at the strain of 0.128, and three instability ranges are exhibited when strain increases to 0.342. It is important to note that the size of instability areas slightly increases with the increasing of strain, which indicates a relevant forming ability of unstable microstructures. Additionally, the effi-
Different regions in the maps indicate specific microstructural evolutions after the specimen were treated by corresponding processing parameters. A narrow unstable region is observed at a lower strain rate range of 0.1–0.3 s1 when the temperature is 800 °C. The corresponding microstructure is shown in Fig. 3(a) and denotes a characteristic of flow instability, which demonstrates that hot working under such conditions should be avoided. It is encouraging that two favorable regions are observed in Fig. 2(b). The first one is in the temperature range of 820–900 °C and a strain rate of 0.1 s1. The peak efficiency is 48%, and it exists at the temperature of 850 °C and the strain rate of 0.1 s1. The microstructure in Fig. 3(b) exhibits a favorable characteristic of DRY. The other favorable region is in the temperature range of 1000–1100 °C and a strain rate range is confined to be less than 0.2 s1. In this region, the peak efficiency is 59% at 1075 °C and the strain rate of 0.1 s1. Accordingly, Fig. 3(c) presents the microstructural features that deformed in this range, in which DRX partially occurred. The microstructural feature of Fig. 3(c) is that some grains are DRX grains (black arrows-necklace structure), and others are deformed grains (other black arrows-pancaked grains). At the lowest strain rate and the highest temperature, the tested steel has almost completely undergone DRX, as seen in Fig. 3(d). Actually, some DRX grains have been proven to grow because of the peak efficiency. This result fits well with current literature [12],
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Fig. 2. Hot processing maps of tested steel at different true strains. (a) e = 0.128; (b) e = 0.342.
Fig. 3. Microstructures of the deformed specimens at different temperatures (arrows in (c) illustrate recrystallized grains and deformed grains). (a) 800 °C, 0.1 s–1; (b) 900 °C, 0.1 s–1; (c) 1000 °C, 0.1 s–1; (d) 1100 °C, 0.1 s–1.
in which DRX occured at the peak efficiency, and was followed by grain coarsening. Hot cracks may be caused in these coarse grains when the material is welded, so it is considered to be unstable. At higher strain rates, every map has two large unstable regions. One is in the lower temperature range, and the other is in the higher temperature range. As strain increases, the unstable areas become more extended and the DRY ranges will be much narrower. Flow instability microstructures such as shear bands tend to be formed at lower temperatures and higher strain rates. In the present investigation, the imposed strain is lower, so no unwanted defects are observed except for common flow instability. Though DRX has almost completely occurred at higher temperature (Fig. 4(b)), the microstructures display a distinct feature of kinked boundaries under this condition (see Fig. 4). The relevant flow stress fluctuates, as demonstrated by its ‘‘wavy’’ behavior,
seen in Fig. 1(c). These kink boundaries are particularly prone to intercrystalline cracking at elevated temperature, and are considered to be unstable. In addition, samples deformed at this elevated temperature range may introduce d-ferrite in uniform austenitic matrix. d-ferrite has deleterious effects on properties, which have been verified in AISI 310 austenitic stainless steel [27]. This result suggests that deformation in this region of the steel should be avoided. Based on a detailed analysis, it can be deduced that the domain at a medium-strain rate and higher temperature is beneficial to the microstructures. Fig. 5(a) [19] shows the microstructure corresponding to the point marked A, which is denoted by a black arrow in Fig. 2(b). Fig. 5(b) and (c) [19] show the microstructures of the samples that deformed at the same temperature and strain rate, but to the larger strains of 0.68 and 1.2, respectively. Fig. 5(b)
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Fig. 4. Microstructures of specimens deformed at different strain rates at 1000 °C and 1100 °C. (a) 1000 °C, e_ ¼ 10 s1; (b) 1100 °C, e_ ¼ 10 s1.
Fig. 5. EBSD images of the samples deformed at 1100 °C and 1 s1 to different strains. (a) e = 0.34 [19]; (b) e = 0.68; (c) e = 1.2 [19].
and (c) show the regular homogeneous austenitic microstructures, and they are different in the degrees of DRX. As seen in Fig. 5(c), DRX varies with the increase of strain, until DRX has been fully occurred. A finer and more uniform microstructure is obtained at the highest strain of 1.2, which supports the claim that this region provides an appropriate forming regime for obtaining homogeneous microstructures. A paradox is discovered between the stress-strain curve of Fig. 1(d) and the microstructure of Fig. 5(b). The former denotes DRY characteristics when the strain is equal to 0.68, but the latter exhibits the trait of DRX. This phenomenon supports the earlier deduction that DRX occurs in the austenitic stainless steel is a
dominant DRY mechanism [20]. Therefore, it is unreasonable to determine whether or not DRX occurred only by the analysis of stress-strain curves. 3.4. Hardness of the deformed sample As shown in Fig. 6, both Vickers hardness and Rockwell hardness of the tested steel gradually reduce with the temperature increase. At the strain rate of 1 s1, the highest hardness is 101HRB, while the lowest one is 76.6 HRB. The reduction percentage is about 24.4%. Similarly, at a strain rate of 0.1 s1, the difference in micro-hardness between the maximum and minimum values is 64 Hv. And the
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Fig. 6. The hardness of tested steel as a function of deformation temperature.
reduction ratio is 23.9%. It is interesting that the degrees of softening measured in the microstructures are roughly equal at these two strain rates. In Fig. 6, both hardness-temperature curves display significant increases near 950 °C. The deviation from the usual value will be larger when the strain rate increased. In order to understand the reason for these changes in hardness, we observed the microstructures of samples deformed at 950 °C. The microstructures show that the increase in hardness is related to the strain induced precipitation (SIP). Momeni et al. [28] studied the competition behavior between SIP and DRX for Fe–28Cr–33Ni–3.5Mo–0.6Mn steel and found that (Cr, Fe, Mo)23C6 precipitates nucleated at the certain temperature and lower strain rate [28]. The precipitates can be observed in the microstructures, in which weak gray GBs and black GBs are etched out, as designated with white and black
arrows in Fig. 7(a), respectively. Detailed examination of the black GBs, as shown in Fig. 7(b), shows that some fine pearl liked precipitates and a spot of white precipitates (indicated with white and black arrows in Fig. 7(b), respectively) are distributed on a high angle GB. Their chemical composition is analyzed by EDS, and the results indicate that they are secondary precipitates of (Cr, Fe, Mo)23C6 and Mo-rich intermetallic compounds of r phase (see Fig. 7(c) and (d)). The compositions in these two kinds of precipitates are very different. There are 5.05C, 73.71Cr, 16.38Fe, 3.87Mo (wt.%) in (Cr, Fe, Mo)23C6 carbides, and 0.28C, 36.31Cr, 44.60Fe, 6.98Mo (wt.%) in sigma precipitates. The precipitation and DRX are diffusion-controlled processes. SIP and DRX have certain inherent connections that may interact and restrict each other. If the temperature is too low, the diffusion rate of solute atoms is low, and precipitation is hard to occur, such as 800 °C. At lower temperatures, recrystallization process is difficult to occur, or may be accompanied by the formation of precipitates, because the solute atoms are liable to cluster at low deformed temperatures and strains [28]. At higher temperature, the diffusion rate of solute atoms is faster. And the pinning of high angle boundaries by SIP is diminished, so DRX is easy to occur. At the highest strain rate, i.e., 10 s1, the time for the occurrence of SIP is too short besides the higher driving force for DRX. Therefore, the DRX is dominated, which resulted in consuming the stored energy, and SIP is restricted. Similarly, during the precipitation, the stored energy in the system is consumed, and the energy dissipation becomes lower. Therefore, the precipitation is related to not only the temperature but also the strain rate. The interaction and restriction between SIP and DRX explain why the energy dissipation percentages near 950 °C are much lower. In this region, the plastic deformation is dominated. And the lager plastic deformation
Fig. 7. OM and SEM observations of precipitates at HAGBs and its EDS spectrum. (a) OM image of sample deformed 950 °C and 0.1 s1; (b) SEM observations of precipitates at HAGBs; (c) EDS spectrum of (Cr, Fe, Mo)23C6 precipitates; (d) EDS spectrum of r phase.
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(4) By the analysis of various regions in the processing maps and the observation of the corresponding microstructures, the workability map of the steel was schematically plotted, which more directly revealed the optimum hot working and instability ranges. The workability map indicated that the optimal deformation parameters of the tested steel were in the temperature range of 1075–1100 °C and the intermediate strain rate range of 0.5–1.7 s1.
Acknowledgements
Fig. 8. Hot workability map of the modified 310 steel under investigated conditions.
increases strain. The strain stimulates the precipitation. As a result, the DRX process is inhibited because of the lowered energy. Based on the analysis, it can be concluded that the precipitation occurred near 950 °C and lower strain rates. 3.5. Hot workability According to the above analysis, the hot workability (processing-microstructural mechanism) map was schematically plotted, which is shown in Fig. 8. For the modified 310 steel, it is obvious that the feasible hot working regime is in the DRX range. Therefore, the optimum hot working parameters of the steel are in the elevated temperature range of 1075–1100 °C and a strain rate range of 0.5–1.7 s1, which will be beneficial for optimizing the microstructures and mechanical properties of the tested steel. It should be noted that the range of appropriate working regimes are rather limited due to the evolution of the material (many coarse carbonization particles on triple joint GBs (Fig. 5, black areas)) and the effect of precipitation during the hot deformation. 4. Conclusions Based on the investigation of the hot processing maps and hot workability of a modified 310 austenitic stainless steel, and the temperature was in a range of 800–1100 °C and strain rate was in a range of 0.1–10 s1, some conclusions were drawn as follows: (1) The energy dissipation rate g and the hot processing maps of the modified 310 steel were obtained based on DMM. The instable deformation regions which should be avoided existed at a high strain rate, and in both low and high temperature ranges. Although the processing maps established at the different strains showed similar features, strain had a significant influence on the processing maps. (2) The deformation in the safe region was beneficial to DRY and DRX, while the deformation in unstable region would lead to flow instability, kink boundaries and grain growth. (3) The hardness of the deformed samples gradually reduced as the temperature increased except that the temperature near 950 °C, at which the hardness evidently increased due to the precipitation hardening. EDS spectroscopy analysis showed that the precipitates were secondary carbonizations of (Cr, Fe, Mo)23C6 and the intermetallic compound of r phase. In addition, because the stored energy was consumed by the precipitation, the SIP resulted in the decrease of energy dissipation rate (g) in the processing maps, and DRX was inhibited.
The authors are grateful for the financial support of Foundation of Science and Technology on Reactor Fuel and Materials Laboratory and the assistance of Xiao Li in discussing the EBSD images in this paper. The authors appreciate Miss. April Novak from University of Illinois for her contribution to the English revision.
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