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Microstructural evolution, precipitation and mechanical properties of hot rolled 27Cr-4Mo-2Ni ferritic steel during 800 °C aging
Materials and Design 160 (2018) 999–1009
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Materials and Design journal homepage: www.elsevier.com/locate/matdes
Microstructural evolution, precipitation and mechanical properties of hot rolled 27Cr-4Mo-2Ni ferritic steel during 800 °C aging Hui-Hu Lu a,b, Hong-Kui Guo a, Yi Luo a, Zhen-Guang Liu c, Wen-Qi Li a, Jian-Chun Li d, Wei Liang a,⁎ a
College of Materials Science and Engineering, Taiyuan University of Technology, Taiyuan 030024, PR China Department of Mechanical and Electrical Engineering, Yuncheng Polytechnic College, Yuncheng 044000, Shanxi Province, PR China College of Materials Science and Engineering, Jiangsu University of Science and Technology, Zhenjiang 212003, Jiangsu Province, PR China d Taiyuan Iron & Steel Co., Ltd., Taiyuan 030003, PR China b c
H I G H L I G H T S
G R A P H I C A L
A B S T R A C T
• Hot-rolling deformation accelerated the formation of intermetallics: chi, Laves and sigma phases. • Formation of intermetallic phases impeded the recovery of deformed ferritin grains. • Precipitation of brittle phases led to brittle fracture mode during tensile test.
a r t i c l e
i n f o
Article history: Received 6 July 2018 Received in revised form 19 October 2018 Accepted 26 October 2018 Available online 28 October 2018 Keywords: Ferritic steel Aging treatment Precipitation Microstructure Embrittlement
a b s t r a c t A 27Cr-4Mo-2Ni super ferritic stainless steel was hot rolled and aged at 800 °C for times ranging from 10 min to 4 h to study the microstructure, precipitation and their effects on mechanical properties by using SEM, EBSD and TEM techniques. Experimental results demonstrate that a great deal of substructure, especially shear bands is obviously observed within elongated ferrite grains after hot rolling. After aging at 800 °C, Laves, chi and sigma phases are identified, and precipitation of these phases is accelerated evidently by the substructure. Laves phase particles are observed formed at dislocations and distributed inside grains with sub-micron scale. Precipitates of chi and sigma phase are found at grain boundaries and along shear bands, as well as around the TiN particles. Recovery during 800 °C aging is impeded by the precipitation of Laves phase at dislocations and chi phase at shear bands. Tensile ductility and micro hardness are significantly affected by the precipitation of intermetallics. The breakage pattern after tensile test transformed from ductile to brittle fracture when the aging time increases from 10 min to 4 h. © 2018 Published by Elsevier Ltd.
1. Introduction
⁎ Corresponding author. E-mail address: [email protected] (W. Liang).
https://doi.org/10.1016/j.matdes.2018.10.039 0264-1275/© 2018 Published by Elsevier Ltd.
Super ferritic stainless steels (SFSS) contain high content of Cr and Mo, and low content of interstitial C and N elements. These steels possess excellent corrosion resistance and high thermal conductivity together with attractive mechanical properties, and they are widely
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China. The chemical composition of the steel was listed in Table 1. These received steels were reheated to 1100 °C for 15 min. Then, they were immediately hot rolled in one pass to a thickness of ~3.1 mm in 10 s and the final temperature of these plates before water quenching was about 1000 °C. Finally, the hot rolled plates were aged treatment at 800 °C for 10 min, 2 h, and 4 h, and then quenched in room temperature water. As a contrast, some received steels were solution treated at 1100 °C for 15 min and immediately quenched in room temperature water. Then, they were subjected to the same aging treatment as hot rolled plates. The schematic graph of processing route is illustrated in Fig. 1. JMatPro software was used to calculate the time-temperatureprecipitate (TTP) curves to forecast the possible intermetallic phases for the tested steel, and the curves are also embedded in Fig. 1. The samples cut from each plate were mechanically ground and then electrochemically polished at 25 V for 30–50 s using the solution consisting of 5 ml perchloric acids and 95 ml absolute ethyl alcohol. The longitudinal section of samples was characterized by field emission scanning electron microscopy (FE-SEM, Tescan Mira 3) equipped with an Oxford energy dispersive spectroscopy (EDS) and an Oxford electron back scattering diffraction (EBSD). The grain size was calculated by EBSD, and at least five hundred grains were determined during the testing. In addition, three positions were detected to calculate the average size and the standard deviation. The bright field images and the selected area diffraction (SAD) pattern of transmission electron microscopy (TEM, JEOL 2100F) were used to show the detailed information of the microstructure. Thin foil samples were prepared electrolytically by double jet electropolishing using a twin jet electropolisher (MTP-1A). Room temperature tensile tests were carried out on a tensile testing machine controlled by electronic signals (DNS 200) with a strain rate of 5 × 10−3 s−1. Sub-size tensile samples were machined with a size of 3.1 × 6 × 25 mm at the gauge region according to the standard of ASTM E8M. The tensile axis was paralleled to the rolling direction (RD). Three interchangeable samples at each condition were used for the tensile testing. The Vickers hardness (HV) was measured with a load of 0.981 N for 15 s on a Vickers hardness tester (HR-320MS). At least six measurements of each condition were conducted for statistical accuracy. The average values and standard deviations were also calculated.
used to manufacture heat exchanger and condenser cooled by corrosive medium, such as heat exchanger for power plants near sea-shore [1–3]. In order to reduce the risk of embrittlement and the degradation of corrosion resistance induced by Cr23C6 or Cr2N, a certain content of Ti and Nb are added into the steels to stabilize the residual C and N by generating TiN and Nb(C, N). Furthermore, 2–4 wt% Ni are usually added into these materials to improve toughness [4,5]. Many kinds of brittle intermetallics are easily formed in the microstructure during thermomechanical treatment processes, such as sigma phase (σ, Fe-Cr or Fe-Cr-Mo), chi phase (χ, Fe36Cr12Mo10) and Laves phase (η, Fe2Nb or Fe2Mo) due to the high content of Cr, Mo and Nb [6–10]. Chi phase belongs to Fe-Cr-Mo ternary system and is characterized by a body centered cubic (BCC, I-43 m space group) structure with a lattice parameter of a = 0.892 nm. Laves phase is approximatively Fe2Mo/Fe2Nb for short with a hexagonal structure (P63/mmc space group) and lattice parameters of a = 0.473 nm and c = 0.772 nm. Sigma phase possesses a tetragonal structure (P42/mmc space group) with lattice parameters of a = 0.88–0.92 nm, c = 0.45–0.48 nm and C = 0.52 nm. Due to the formation of these intermetallics, brittle cracks or fracture phenomena appear frequently during cold rolling or unfolding process in the industrial production of SFSSs [6,7,11]. Thus, high temperature solution treatment and rapid cooling are necessary to recover the mechanical properties of SFSSs before cold rolling [12,13]. It should be noted that the thermomechanical process is required if the brittle precipitates are formed during inappropriate process, because simple solution treatment at high temperature easily induces unexpected grains coarsening which is detrimental to mechanical properties [2]. Formation of intermetallic phases, chi, Laves and sigma, often occurs in the temperature range from 600 to 1000 °C in several kinds of SFSSs, such as 25Cr-3Mo-4Ni [8], 26Cr-4Mo–2Ni [14], 28Cr-4Mo [6,7], 28Cr4Ni-2Mo [15], 29Cr-4Mo-2Ni [16–18], 30Cr-4Mo-2Ni [13] and 39Cr2Mo-2Ni [5]. The kinetics curves of these precipitation display a “C” shape, and the “nose” temperature of all the temperature-time precipitation (TTP) curves located between 800 and 900 °C [8]. When these alloys are aged at intermediate temperatures, chi phase particles usually locate at the grain boundaries and form networks of grain boundary with prolonged aging. Sigma precipitates are observed in association with chi phase particles and formed at grain boundaries [8,14,17,and]. Sigma phase is the dominant constituent in samples produced at the longer aging times and possesses a dendritic morphology [8]. Laves phase precipitates not only at grain boundaries [15] but also randomly through the microstructure [14]. To date, many studies were carried out to study the precipitation behavior of the steels during isothermal aging after solution treatment. However, little attention has been paid to the subject of the effect of hot-rolling deformation on the evolution of microstructure, precipitation and mechanical properties of 27Cr-4Mo-2Ni SFSSs during intermediate temperature aging, especially near the “nose” temperature. In this paper, 27Cr-4Mo-2Ni SFSSs subjected to hot rolling and water quenching are aged at 800 °C. Meanwhile, the effect of hot-rolling deformation on microstructure, precipitation, tensile properties and micro hardness are researched aiming to understand of the relationship between microstructure and mechanical properties of SFSSs.
3. Results and discussion 3.1. Microstructure and tensile properties of as-hot rolled steels Fig. 2 shows the microstructure and precipitates of samples after solution treating and hot rolling. Due to the high content of Cr and Mo together with low content of C and N in the steel, fully recrystallized ferrite grains are observed after solution treating (Fig. 2a). It is measured that the average size of ferrite grains is about 69 ± 3 μm. The results of EDS analysis reveal that the black square and light gray phases are TiN and Nb(C, N) particles, respectively (Fig. 2e and f). These two kinds of phases are formed during the stabilization reaction of C or N and Ti or Nb at high temperature, and they are thermodynamically stable during the medium temperature heat treatment [13,14]. The Nb(C, N) particles are particularly distributed along RD in a line due to the inheritance from the initial industrial hot rolled plates [12], in which the Nb(C, N) particles locate at the initial grain boundaries along RD (Fig. 2b). After hot rolling, a deformed microstructure undergoing dynamic recovery by means of slip and climb of dislocations was produced due to high stacking fault energy in SFSSs [19]. The ferritic grains are slightly elongated along the rolling direction and contain a great deal of
2. Experimental procedures A 27Cr-4Mo-2Ni SFSS stabilized by Nb and Ti was continuous cast into slabs with a thickness of ~200 mm. These cast slabs were reheated to 1200 °C and then hot rolled into 4.1 mm with passes in TISCO of Table 1 The chemical composition of 27Cr-4Mo-2Ni SFSS (wt%). C
Si
Mn
P
S
Cr
Ni
Mo
Cu
Nb
Ti
N
Fe
0.015
0.4
0.23
0.022
0.002
27.57
1.98
3.72
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Fig. 1. Schematic graph of fabrication procedure and TTP-curves for the 27Cr-4Mo-2Ni alloy.
substructure characterized by low angle grain boundaries (~88.1%), which are induced by the weak hot deformation [20], as shown in Fig. 2c and the following Fig. 5a. The deformation in different grains is heterogeneous with an average grain size of 65 ± 3 μm, as shown in Fig. 2g. Some shear bands inside the heavily deformed grains characterized by {111} orientation are also observed. This finding agrees with the reported deformation texture in ferritic steels [21]. Whereas the TiN and Nb(C, N) particles are kept relatively steady in the microstructure (Fig. 2d). Fig. 3 shows the typical engineering stress-strain curves of samples processed by hot rolling and solution treating. After solution treating at 1100 °C for 15 min, the ultimate tensile strength (UTS) of the sample is as high as 690 ± 11 MPa and the break elongation (EL) is 20 ± 2% due
3.2. Microstructural evolution of aged samples Fig. 4 shows the microstructural evolution of hot rolled samples during isothermal aging. It is obvious that a great deal of low angle grain boundaries are observed and recrystallization process is not activated, because the completed recrystallization temperature of hot rolled SFFSs is N1050 °C [12,13]. After 800 °C aging, the deformed microstructure characterized by dynamic recovery are reserved together with a further static recovery, which is also confirmed by many grain boundaries with misorientation angle of 2–5° inside the grains, as shown in Fig. 5. Nevertheless, the static recovery kinetics during the isothermal
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Fig. 2. EBSD images (a and c) and BSE (back scattered electron) images (b and d) showing the microstructure and precipitates of (a and b) the solution annealed and (c and d) the hot rolled plate, respectively. The SEM-EDS of TiN and Nb(C,N) are shown in (e) and (f). The grain size distribution of panel c is shown in panel (g). In (a) and (c), grain boundaries with misorientation larger than 15° are colored as black lines and the white lines represent low-angle grain boundaries with misorientation of 2–15°.
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(a)
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200μm Fig. 4. EBSD images showing the microstructure of the hot rolled samples after 800 °C aging treatment for different times: (a) 10 min, (b) 2 h and (c) 4 h. Grain boundaries with misorientation larger than 15° are colored as black lines. The white lines represent low-angle grain boundaries with misorientation of 2–15°.
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Fig. 5. Misorientation between grains in samples after (a) hot rolling and aging treatment at 800 °C for (b) 10 min, (c) 2 h and (d) 4 h. The distribution of low angle grain boundaries (b15°) is inserted in each image.
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aging is sluggish. For example, the fraction of low angle grain boundaries (b5°) in the sample aged for 10 min is about 75.0%, whereas the values for 2 h and 4 h increase slightly to ~76.9% and ~78.3%, respectively (Fig. 5). The average grain size of samples aged for 10 min, 2 h and 4 h, is about 65 ± 3, 67 ± 3 and 67 ± 4 μm, respectively. It should be noted that the growth of grains is not observed with increasing aging time, which is different from the results of solution treated (ST) samples. The grain size of aged samples without hot rolling increases from 67 ± 3 (ST) to 110 ± 7 μm (aged for 4 h). This contributes to the inhibition effect of precipitates on grain boundaries [14]. The precipitation of intermetallics is discussed in Section 3.3. Fig. 6 shows the microtexture of these samples suffering from different thermomechanical treatment. After hot rolling with reduction rate of ~25%, typical α-fiber (i.e., 〈110〉//RD, rolling direction) and γ-fiber (i.e., 〈111〉//ND, normal direction) texture are developed [23] and the
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peak value appears near {111} 〈123〉 component. After aging at 800 °C, the γ-fiber texture is weakened, while the α-fiber texture is strengthened. With increasing aging time, grain orientation is gradually transformed from γ-fiber texture to α fiber texture and the peak value evolves near {115} 〈110〉 component. 3.3. Effect of hot rolling on precipitation evolution Fig. 7 shows the TEM images and SAD of typical intermetallics in the samples aged at 800 °C. The results show that three kinds of intermetallics, Laves, chi, and sigma phase, are identified by TEM, which is in accord with the results predicted by JMatPro, as shown in the TTP curves embedded in Fig. 1. It is obvious from the TTP curves that precipitation of intermetallics phases, Laves, chi, and sigma, occur in the temperature
Fig. 6. Constant φ2 = 45° ODF(oriented distribution function) images showing the microtexture of samples after different thermomechanical processing: (a) solution treatment, (b) hot rolling, (c) aging at 800 °C for 10 min, (d) aging at 800 °C for 2 h and (e) aging at 800 °C for 4 h.
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Fig. 7. TEM images and SAD pattern of (a and b) Laves, (c and d) chi and (e and f) sigma phases in the hot rolled samples aged at 800 °C for (a and c) 10 min and (e) 2 h.
range 600–1000 °C. The kinetics curves of these precipitation display a C-shape with the “nose” temperature located between 800 and 950 °C. The TTP curves also show that the nucleation of Laves phase is faster than chi or sigma phase, and the formation temperature of Laves and chi phases is higher than that of sigma phase. The “nose” temperatures located between 800 and 950 °C of the calculated TTP curves in this research are higher than that in a 25Cr-3Mo-4Ni steel [8]. This is because that the high content of Cr and Mo enhances the stability of chi, sigma and Laves phases [5–8,13,14]. Whereas, the “nose” temperatures of these phases in 25Cr-3Mo-4Ni steels are uniform, which differs from the data in this paper. This contributes to the difference of chemical composition between 25Cr-3Mo-4Ni and 27Cr-4Mo-2Ni steels in this study. The morphology and distribution of intermetallic phases as a function of aging time are illustrated in Fig. 8. In Figs. 7 and 8, the rectangular Laves phase nucleates originally at dislocations and dislocation walls (Fig. 7a), and then develops into needle-like particles with the size of 0.1–0.3 μm. They are dispersedly distributed through the microstructure with increasing aging time (Fig. 8e). In the early stage of isothermal aging, the precipitation of fine chi phase appears at grain boundaries (Figs. 7c and 8a). With an increase in aging time, the amount of chi phase particles increases constantly and they form continuous networks along grain boundaries, as shown in Fig. 8b and c. After aging for 4 h, chi phase particles are found both along shear bands and around the TiN particles (Fig. 8c, d and e). Shear bands are marked within yellow dotted ellipses in Fig. 8c. Bulk sigma phase particles were firstly found around chi precipitates along grain boundaries. After aging for 2 h, sigma phase particles occupy most grain boundaries (Fig. 8b) and expand
into the interior of grains with a dendritic morphology (marked by red line with an arrow in Fig. 8d). The TEM image in Fig. 7d also shows the dendritic morphology of sigma phase. As the aging time increases to 4 h, sigma phase precipitates are observed around chi phase particles along shear bands (marked by blue line with an arrow in Fig. 8c and d). In addition, a few bulk sigma particles are also found to be distributed in the vicinity of TiN particles and some isolated sigma phase particles are also observed in the interior of grains (Fig. 8c). It should be pointed out that a number of microcracks are also observed (Fig. 8d) due to the result of the extremely brittle nature of this phase [5,14]. The results of SEM-EDS analysis in Table 2 show that the elements of chi phase are same as that of sigma phase: Fe, Cr, Mo Si and Ni. However, chi phase contains more Mo and less Cr than that of sigma phase. Laves phase is composed of Fe, Cr, Mo, Nb, Si, and Ni, and is enriched in Nb compared with both sigma and chi phases. One major objective of this study is to clarify the effect of hot rolling on precipitation reactions that occur in a 27Cr-4Mo-2Ni SFSS. Fig. 9 displays the microstructure of solution treated samples after isothermal aging for comparing. The results show that chi phase only forms at grain boundaries (Fig. 9a and b). A few sigma phase particles are observed and randomly distributed inside the grains except for the main precipitation of sigma phase at grain boundaries (Fig. 9b). Laves phases are distributed in the interior of grains (Fig. 9a and b). It is well known that second phases in steels preferentially precipitate at crystal defects, such as the boundaries of grain, subgrain and particles of other phases, dislocations, and vacancies [22]. In the present study, chi phase precipitates in samples aged for 10 min appears at grain boundaries due to the low barrier for heterogeneous nucleation. With increasing aging time,
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Fig. 8. BSE images showing the morphology and distribution of typical intermetallic in the hot rolled samples after aging treatment at 800 °C for different times: (a) 10 min, (b) 2 h, (c, d and e) 4 h. The image (d) and (e) are the high magnification picture of the ellipse area in the image (c).
the diffusion of Cr and Mo atoms, which compose the main elements of chi phase is sufficient [17,24]. Thus, shear bands characterized by low angle boundaries provide extra sites for the heterogeneous nucleation of chi phase (Fig. 8e) due to the high energy in shear bands, which offer the necessary drive force for the precipitation of chi phase [25,26]. As a result, the chi phase precipitated at both grain boundaries and shear bands are detected during long time aging. Precipitation of
sigma phase preferentially appears around the chi phase particles at boundaries, because areas surrounding chi phase particles, such as grain boundaries and shear bands possess enough Mo and Cr atoms [17,27,28]. Stress concentration around TiN particles is produced during hot rolling processes because of the different deformation ability of the ferrite matrix and TiN particles. Thus, high interfacial energy is generated, which is in favour of the formation of chi phase and sigma phase.
Table 2 EDS results of typical precipitates in 27Cr-4Mo-2Ni SFSS aged at 800 °C (wt%). Elements
Fe
Cr
Mo
Nb
Si
Ni
Matrix Sigma phase Chi phase Laves phase
67.1 ± 1.1 61.3 ± 1.5 58.0 ± 0.8 53.6 ± 0.8
27.8 ± 1.3 31.8 ± 1.2 22.8 ± 0.9 20.2 ± 0.8
3.0 ± 0.4 4.9 ± 0.5 15.7 ± 1.0 10.6 ± 0.9
0.3 ± 0.1 – – 10.3 ± 0.8
0.1 ± 0.1 0.4 ± 0.3 1.4 ± 0.4 1.5 ± 0.4
1.7 ± 0.1 1.6 ± 0.4 2.1 ± 0.2 3.8 ± 1.1
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Fig. 9. BSE images showing the morphology and distribution of typical intermetallic in the solution treated samples after aging treatment at 800 °C for (a) 2 h and (b) 4 h.
Moreover, the area fraction of the sigma phase is calculated by Image Pro, as shown in Fig. 10. It is obvious that sigma phase precipitation is accelerated by hot-rolling deformation, especially at long time aging. The area fraction is increased from 5.3 ± 0.4% to 8.7 ± 0.7% when aged for 4 h at 800 °C. The dislocations induced by hot rolling are suitable nucleation sites for Laves precipitation, as shown in Fig. 7a. This phenomenon is also found in other high Cr ferritic steels [8,29]. The dislocations in the crystal provide a path for the diffusion of solution atoms, such as Nb in this study, which promotes the diffusion reactions [30–32]. In consequence, a large quantity of Laves phase is observed at dislocations and dislocation cells. The size of the Laves phase in the hot rolled samples after aging for 4 h is 0.1–0.3 μm. While the size of Laves phase in the samples without hot rolling is 0.5–1.5 μm, which is similar to that in other report [14]. This difference was because that the dislocations and substructure produced by hot-rolling deformation provided additional sites for the nucleation of Laves phase particles. 3.4. Effect of precipitation on microstructural evolution Many studies show that there is an interaction between microstructure and precipitation [33]. The second phases at the grain boundaries or subgrain boundaries impede the recovery and recrystallization of ferrite grains during annealing processes [33–35]. In the present study, recovery of the deformed microstructure is dominant during the 800 °C aging, as shown in Fig. 5. After aging for 10 min at 800 °C, formation of chi phase at grain boundaries and Laves phase at dislocation walls are observed (Figs. 7a and 8a). The migration of grain boundaries during the further long time aging is inhibited due to the pinning effect of fine chi phase particles [14,29]. It is obvious that the grain boundaries
are completely covered by the chi phase (Figs. 8 and 9). Besides, due to the presence of intermetallic phases at grain boundaries and the relative high aging temperature (800 °C), solute drag effects may provide an additional impact [29]. Thus, the average grain sizes of hot rolled samples after aging are nearly constant. In addition, the recovery kinetics is sluggish due to the inhibition of fine Laves phase precipitates on the movement of dislocations, as shown in Fig. 11a and b. The Laves phase particles are marked by a yellow arrow in Fig. 11a. In general, the deformation in the microstructure of ferritic steels is heterogeneous due to different deformability between α-fiber and γfiber grains, and a few shear bands are usually observed in the γ-fiber grains [34]. In addition, strong recovery of the α-fiber grains are usually observed due to the kinetic and thermodynamic stability after recovery annealing [35]. In the present study, many fine chi phase particles are found along shear bands in γ-fiber grains (Figs. 4b and 8e). During medium temperature aging, such as 800 °C in this study, it is possible to further initiate recovery for α-fiber grains by way of slip and climb of dislocations despite of the hindrance of fine Laves phase particles on the movement of dislocations [14]. The subgrains are clearly observed in the TEM image of samples aged for 10 min, and the inner section of subgrain is relatively smooth, as shown in Fig. 11a. In contrast, the development of γ-fiber grains is weak due to the additional pinning effect of fine chi phase precipitates on shear bands although their storage energy is high [36,37]. The kernel average misorientation (KAM) in Fig. 11d also shows the higher energy in shear bands (higher KAM values). Therefore, recovery grains are rarely observed at shear bands (marked by black arrows), as shown in Fig. 11b and c. As a result, the α-fiber texture is strengthened slightly, while the γ-fiber texture is weakened with increasing aging time, as shown in Fig. 6. This finding agrees with the results in a FeCrAl stainless steel reported by Z. Sun [34].
Area fraction of sigma phase/%
10
3.5. Effect of microstructure and precipitation on microhardness
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Fig. 10. Histogram showing the effect of hot rolling on area fraction of sigma phase after 800 °C aging.
For Vickers hardness testing, the average value of each specimen is calculated and the results are shown in Fig. 12. It is obvious that the microhardness of hot rolled samples is higher than that of solution treated ones due to the work hardening induce by hot rolling. The microhardness of aged samples decreases compared with the hot rolled one, and then increases continuously with increasing aging time. Both the reduction of work hardening induced by the additional aging recovery and the loss of solution strengthening caused by the precipitation of intermetallics account for the slight decrease of the hardness of samples aged for 10 min. The continuous increasing of hardness with increasing aging time is mainly due to the fast increase of both the size and content of sigma precipitates [12,14], as shown in Fig. 8 that the breadth of sigma phase particles extends to 15 μm, and in Fig. 10 that the area fraction of sigma phase particles increases to 8.7 ± 0.7%. Besides, the Vickers
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Fig. 11. TEM image (a), BC (band contrast) map (b), IPF (inverse pole figure) colored map (c) and KAM (kernel average misorientation) map (d) showing the recovery in samples aged at 800 °C for (a) 10 min, and (b, c and d) 2 h. The white lines in image (c) represent low-angle grain boundaries with misorientation of 2–15°.
hardness of hot rolled samples after aging is much higher than that of these solution treated ones. In general, the Vickers hardness of ferritic steels is related to the area fraction of precipitates [12,14]. Thus, the larger number of precipitates account for the increase in value of hardness for the hot rolled samples. The main reason is that the precipitation of intermetallics is evidently promoted by the hot-rolling deformation in the microstructure, as shown in Fig. 10. The slope of the curves of the sample aged for 4 h is larger than that for 2 h, indicating that the increase of Vickers hardness is accelerated by long aging. This is because that many additional chi and sigma phases form along shear bands in samples aged for 4 h (Fig. 10). Especially, the slope of the hardness curve between 10 min and 2 h in hot rolled samples is less than that in aged samples without hot rolling. Residual work hardening in the 400 Hot rolled +Aged Solution treated + Aged
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Thermomechanical treatment Fig. 12. Vickers hardness of samples with different thermomechanical treatments. (ST: solution treated, HR: hot rolled).
sample with 10 min aging after hot rolling makes the value of hardness rise to 295 ± 11 HV and then leads to a relatively gentle slope compared with the ones aged without hot rolling. 3.6. Effect of microstructure and precipitation on mechanical properties Fig. 13 shows the comparison of tensile properties for samples with different thermomechanical treatment. It is obvious that both the tensile strength and elongation (EL) of the samples decrease dramatically with increasing aging time. While the tensile strength of aged samples with hot rolling is lower than that of the solution treated ones, as shown in Fig. 13a. After aging for 10 min, the EL of hot rolled samples is about 20 ± 2%, while it decreases to 4 ± 0.8% at a long aging time for 4 h. The EL of aged samples without hot rolling are similar to the values of the hot rolled ones, as shown in Fig. 13b. In general, the strengthening mechanisms of ferritic steels mainly include solution strengthening, work hardening and precipitation strengthening [13,22]. Compared with hot rolled samples, the tensile strength decreases with a reduction of ~120 MPa due to the loss of work hardening induced by the process of recovery when the sample is aged at 800 °C for 10 min. However, the value of EL is improved to the grade of the samples treated by solution, because limited chi phase particles are observed in the grain boundaries. With an increase in aging time, the tensile strength of hot rolled samples aged for 2 h decreases to 583 ± 8 MPa and further to 529 ± 8 MPa after 4 h aging. This is because that precipitation and coarsening of intermetallics consume plenty of Cr, Mo and Si atoms, which lead to heavy loss of the solution strengthening [13,14]. In addition, the higher loss of the solution strengthening induced by more precipitation of intermetallics also accounts for the decrease of the tensile strength of hot rolled samples after aging compared to the solution treated ones. The precipitation of intermetallics at grain boundaries,
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(b) 800
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Fig. 13. Comparison of tensile properties for samples with different thermomechanical treatments: (a) tensile strength and (b) elongation. (ST: solution treated, HT: hot rolled).
such as chi and sigma phases accounts for the rapid decline of EL. During tensile testing, the samples aged for 4 h suffer from a sudden fracture because most grain boundaries are covered by brittle chi and sigma phases. Besides, the microcracks throughout the sigma phase present an additional pernicious effect on the deterioration of EL, as shown in Fig. 8d. Accordingly, the fracture surface of the samples after tensile testing is observed by SEM, and the results are exhibited in Fig. 14. Many dimples are detected in the fracture surface of the samples aged for 10 min (Fig. 14a), which indicates a ductile fracture mode. In contrast, river pattern and cleavage facets are predominant on the surface of samples aged for 2–4 h (Fig. 14b and c). This presents a typical mode of brittle fracture [12,38]. It should be noted that precipitation and coarsening of grain boundary phases lead to a sharp decline of EL and make the
breakage pattern transformed from ductile to brittle fracture when the aging time increases from 10 min to 4 h. 4. Conclusions A 27Cr-4Mo-2Ni super ferritic stainless steel was hot rolled and subjected to isothermal aging at 800 °C for times ranging from between 10 min and 4 h. The evolution of microstructure and precipitation, and their effects on mechanical properties are investigated, and a few conclusions are drawn as follows: 1. After hot rolling, recovered ferrite grains with many low angle grain boundaries and shear bands, and embedded TiN and Nb(C,
Fig. 14. SEM images showing the impact fracture morphology of tested steels after aging treatment at 800 °C for different time: (a) 10 min, (b) 2 h and (c) 4 h.
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N) particles are observed in the microstructure. The tensile strength and elongation are 731 ± 10 MPa and 14 ± 3% respectively. 2. After aging at 800 °C, Laves phase nucleates at the location of dislocations and is distributed inside the grains. The nanosized chi phase particles start to precipitate at grain boundaries after short time aging. These chi phases form networks along grain boundaries and nucleate along shear bands with prolonged aging time. The precipitation of sigma phase appears around chi phase particles which locate at grain boundaries, shear bands and around TiN particles. Sigma phase is the dominant constituent and shows a dendritic morphology with increasing aging time. 3. Precipitation of intermetallics is accelerated by hot-rolling deformation. Substructure induced by hot rolling provides additional nucleation sites for intermetallics precipitation including chi and sigma phases. Plenty of dislocations in recovered grains present preferential locations for Laves phase precipitation. Recovery of hot rolled plate during aging is impeded by the formation of chi and Laves phases at dislocation and shear bands, respectively. 4. Vickers hardness increases with increasing aging time due to the higher content of the intermetallic phases formed after longer time aging. The value of hardness of hot rolled samples after aging is higher than that of aged samples without hot rolling. The tensile strength and elongation decrease dramatically with prolonged aging. Precipitation and coarsening of grain boundary phases account for the sharp decline of elongation and makes the breakage pattern transformed from ductile to brittle fracture when the aging time increases from 10 min to 4 h. Author contribution statement The first author (Hui-Hu Lu) presented the research idea and experimental design, and wrote this manuscript. The second and third authors (Hong-Kui Guo and Yi Luo) assisted to complete the hot rolling and annealing experiment. The fourth author (Zhen-Guang Liu) put forward some instructive suggestions for the writing of this paper and modified the language. The fifth author (Wen-Qi Li) helped to fulfill the work of SEM and EBSD analysis. The sixth author (Jian-Chun Li) helped to carry out rolling experiment and calculate the TTP curves. The TEM analysis was completed by the corresponding author (Wei Liang), and he was the faculty adviser of the first author. Acknowledgements This work was supported by Projects of International Cooperation in Shanxi (Grant No. 201603D421026). References [1] D. Janikowski, W. Henricks, Super-ferritic Stainless Steels - The Cost Effective Answer for Heat Transfer Tubing, REF: MP07-016, World Congress on Desalination and Water Reuse, Gran Canarias, Spain, 2007 (October 21–26). [2] P.A. Olubamb, J.H. Potgieter, L. Cornish, Corrosion behaviour of superferritic stainless steels cathodically modified with minor additions of ruthenium in sulphuric and hydrochloric acids, Mater. Des. 30 (2009) 1451–1457. [3] N. Pessall, J.I. Nurminen, Development of ferritic stainless steels for use in desalination plants, Corrosion 30 (11) (2013) 381–392. [4] I.A. Franson, Mechanical properties of high purity Fe-26 Cr-1 Mo ferritic stainless steel, Metall. Trans. 5 (1974) 2257–2264. [5] K. Premachandra, M.B. Cartie, R.H. Eric, Effect of stabilising elements on formation of σ phase in experimental ferritic stainless steels containing 39%Cr, Mater. Sci. Technol. 8 (2013) 437–442. [6] M.A. Streicher, Microstructures and some properties of Fe-28%Cr-4%Mo alloys, Corrosion 30 (4) (1974) 115–124. [7] T.J. Nichol, Microstructures and some properties of Fe-28%Cr-4%Mo alloys, Metall. Trans. A. 8 (1977) 229–237. [8] E.L. Brown, M.E. Burnett, P.T. Purtscher, G. Krauss, Intermetallic phase formation in 25Cr-3Mo-4Ni ferritic stainless steel, Metall. Trans. A. 14 (1983) 791–800.
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Materials and Design journal homepage: www.elsevier.com/locate/matdes
Microstructural evolution, precipitation and mechanical properties of hot rolled 27Cr-4Mo-2Ni ferritic steel during 800 °C aging Hui-Hu Lu a,b, Hong-Kui Guo a, Yi Luo a, Zhen-Guang Liu c, Wen-Qi Li a, Jian-Chun Li d, Wei Liang a,⁎ a
College of Materials Science and Engineering, Taiyuan University of Technology, Taiyuan 030024, PR China Department of Mechanical and Electrical Engineering, Yuncheng Polytechnic College, Yuncheng 044000, Shanxi Province, PR China College of Materials Science and Engineering, Jiangsu University of Science and Technology, Zhenjiang 212003, Jiangsu Province, PR China d Taiyuan Iron & Steel Co., Ltd., Taiyuan 030003, PR China b c
H I G H L I G H T S
G R A P H I C A L
A B S T R A C T
• Hot-rolling deformation accelerated the formation of intermetallics: chi, Laves and sigma phases. • Formation of intermetallic phases impeded the recovery of deformed ferritin grains. • Precipitation of brittle phases led to brittle fracture mode during tensile test.
a r t i c l e
i n f o
Article history: Received 6 July 2018 Received in revised form 19 October 2018 Accepted 26 October 2018 Available online 28 October 2018 Keywords: Ferritic steel Aging treatment Precipitation Microstructure Embrittlement
a b s t r a c t A 27Cr-4Mo-2Ni super ferritic stainless steel was hot rolled and aged at 800 °C for times ranging from 10 min to 4 h to study the microstructure, precipitation and their effects on mechanical properties by using SEM, EBSD and TEM techniques. Experimental results demonstrate that a great deal of substructure, especially shear bands is obviously observed within elongated ferrite grains after hot rolling. After aging at 800 °C, Laves, chi and sigma phases are identified, and precipitation of these phases is accelerated evidently by the substructure. Laves phase particles are observed formed at dislocations and distributed inside grains with sub-micron scale. Precipitates of chi and sigma phase are found at grain boundaries and along shear bands, as well as around the TiN particles. Recovery during 800 °C aging is impeded by the precipitation of Laves phase at dislocations and chi phase at shear bands. Tensile ductility and micro hardness are significantly affected by the precipitation of intermetallics. The breakage pattern after tensile test transformed from ductile to brittle fracture when the aging time increases from 10 min to 4 h. © 2018 Published by Elsevier Ltd.
1. Introduction
⁎ Corresponding author. E-mail address: [email protected] (W. Liang).
https://doi.org/10.1016/j.matdes.2018.10.039 0264-1275/© 2018 Published by Elsevier Ltd.
Super ferritic stainless steels (SFSS) contain high content of Cr and Mo, and low content of interstitial C and N elements. These steels possess excellent corrosion resistance and high thermal conductivity together with attractive mechanical properties, and they are widely
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China. The chemical composition of the steel was listed in Table 1. These received steels were reheated to 1100 °C for 15 min. Then, they were immediately hot rolled in one pass to a thickness of ~3.1 mm in 10 s and the final temperature of these plates before water quenching was about 1000 °C. Finally, the hot rolled plates were aged treatment at 800 °C for 10 min, 2 h, and 4 h, and then quenched in room temperature water. As a contrast, some received steels were solution treated at 1100 °C for 15 min and immediately quenched in room temperature water. Then, they were subjected to the same aging treatment as hot rolled plates. The schematic graph of processing route is illustrated in Fig. 1. JMatPro software was used to calculate the time-temperatureprecipitate (TTP) curves to forecast the possible intermetallic phases for the tested steel, and the curves are also embedded in Fig. 1. The samples cut from each plate were mechanically ground and then electrochemically polished at 25 V for 30–50 s using the solution consisting of 5 ml perchloric acids and 95 ml absolute ethyl alcohol. The longitudinal section of samples was characterized by field emission scanning electron microscopy (FE-SEM, Tescan Mira 3) equipped with an Oxford energy dispersive spectroscopy (EDS) and an Oxford electron back scattering diffraction (EBSD). The grain size was calculated by EBSD, and at least five hundred grains were determined during the testing. In addition, three positions were detected to calculate the average size and the standard deviation. The bright field images and the selected area diffraction (SAD) pattern of transmission electron microscopy (TEM, JEOL 2100F) were used to show the detailed information of the microstructure. Thin foil samples were prepared electrolytically by double jet electropolishing using a twin jet electropolisher (MTP-1A). Room temperature tensile tests were carried out on a tensile testing machine controlled by electronic signals (DNS 200) with a strain rate of 5 × 10−3 s−1. Sub-size tensile samples were machined with a size of 3.1 × 6 × 25 mm at the gauge region according to the standard of ASTM E8M. The tensile axis was paralleled to the rolling direction (RD). Three interchangeable samples at each condition were used for the tensile testing. The Vickers hardness (HV) was measured with a load of 0.981 N for 15 s on a Vickers hardness tester (HR-320MS). At least six measurements of each condition were conducted for statistical accuracy. The average values and standard deviations were also calculated.
used to manufacture heat exchanger and condenser cooled by corrosive medium, such as heat exchanger for power plants near sea-shore [1–3]. In order to reduce the risk of embrittlement and the degradation of corrosion resistance induced by Cr23C6 or Cr2N, a certain content of Ti and Nb are added into the steels to stabilize the residual C and N by generating TiN and Nb(C, N). Furthermore, 2–4 wt% Ni are usually added into these materials to improve toughness [4,5]. Many kinds of brittle intermetallics are easily formed in the microstructure during thermomechanical treatment processes, such as sigma phase (σ, Fe-Cr or Fe-Cr-Mo), chi phase (χ, Fe36Cr12Mo10) and Laves phase (η, Fe2Nb or Fe2Mo) due to the high content of Cr, Mo and Nb [6–10]. Chi phase belongs to Fe-Cr-Mo ternary system and is characterized by a body centered cubic (BCC, I-43 m space group) structure with a lattice parameter of a = 0.892 nm. Laves phase is approximatively Fe2Mo/Fe2Nb for short with a hexagonal structure (P63/mmc space group) and lattice parameters of a = 0.473 nm and c = 0.772 nm. Sigma phase possesses a tetragonal structure (P42/mmc space group) with lattice parameters of a = 0.88–0.92 nm, c = 0.45–0.48 nm and C = 0.52 nm. Due to the formation of these intermetallics, brittle cracks or fracture phenomena appear frequently during cold rolling or unfolding process in the industrial production of SFSSs [6,7,11]. Thus, high temperature solution treatment and rapid cooling are necessary to recover the mechanical properties of SFSSs before cold rolling [12,13]. It should be noted that the thermomechanical process is required if the brittle precipitates are formed during inappropriate process, because simple solution treatment at high temperature easily induces unexpected grains coarsening which is detrimental to mechanical properties [2]. Formation of intermetallic phases, chi, Laves and sigma, often occurs in the temperature range from 600 to 1000 °C in several kinds of SFSSs, such as 25Cr-3Mo-4Ni [8], 26Cr-4Mo–2Ni [14], 28Cr-4Mo [6,7], 28Cr4Ni-2Mo [15], 29Cr-4Mo-2Ni [16–18], 30Cr-4Mo-2Ni [13] and 39Cr2Mo-2Ni [5]. The kinetics curves of these precipitation display a “C” shape, and the “nose” temperature of all the temperature-time precipitation (TTP) curves located between 800 and 900 °C [8]. When these alloys are aged at intermediate temperatures, chi phase particles usually locate at the grain boundaries and form networks of grain boundary with prolonged aging. Sigma precipitates are observed in association with chi phase particles and formed at grain boundaries [8,14,17,and]. Sigma phase is the dominant constituent in samples produced at the longer aging times and possesses a dendritic morphology [8]. Laves phase precipitates not only at grain boundaries [15] but also randomly through the microstructure [14]. To date, many studies were carried out to study the precipitation behavior of the steels during isothermal aging after solution treatment. However, little attention has been paid to the subject of the effect of hot-rolling deformation on the evolution of microstructure, precipitation and mechanical properties of 27Cr-4Mo-2Ni SFSSs during intermediate temperature aging, especially near the “nose” temperature. In this paper, 27Cr-4Mo-2Ni SFSSs subjected to hot rolling and water quenching are aged at 800 °C. Meanwhile, the effect of hot-rolling deformation on microstructure, precipitation, tensile properties and micro hardness are researched aiming to understand of the relationship between microstructure and mechanical properties of SFSSs.
3. Results and discussion 3.1. Microstructure and tensile properties of as-hot rolled steels Fig. 2 shows the microstructure and precipitates of samples after solution treating and hot rolling. Due to the high content of Cr and Mo together with low content of C and N in the steel, fully recrystallized ferrite grains are observed after solution treating (Fig. 2a). It is measured that the average size of ferrite grains is about 69 ± 3 μm. The results of EDS analysis reveal that the black square and light gray phases are TiN and Nb(C, N) particles, respectively (Fig. 2e and f). These two kinds of phases are formed during the stabilization reaction of C or N and Ti or Nb at high temperature, and they are thermodynamically stable during the medium temperature heat treatment [13,14]. The Nb(C, N) particles are particularly distributed along RD in a line due to the inheritance from the initial industrial hot rolled plates [12], in which the Nb(C, N) particles locate at the initial grain boundaries along RD (Fig. 2b). After hot rolling, a deformed microstructure undergoing dynamic recovery by means of slip and climb of dislocations was produced due to high stacking fault energy in SFSSs [19]. The ferritic grains are slightly elongated along the rolling direction and contain a great deal of
2. Experimental procedures A 27Cr-4Mo-2Ni SFSS stabilized by Nb and Ti was continuous cast into slabs with a thickness of ~200 mm. These cast slabs were reheated to 1200 °C and then hot rolled into 4.1 mm with passes in TISCO of Table 1 The chemical composition of 27Cr-4Mo-2Ni SFSS (wt%). C
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Fig. 1. Schematic graph of fabrication procedure and TTP-curves for the 27Cr-4Mo-2Ni alloy.
substructure characterized by low angle grain boundaries (~88.1%), which are induced by the weak hot deformation [20], as shown in Fig. 2c and the following Fig. 5a. The deformation in different grains is heterogeneous with an average grain size of 65 ± 3 μm, as shown in Fig. 2g. Some shear bands inside the heavily deformed grains characterized by {111} orientation are also observed. This finding agrees with the reported deformation texture in ferritic steels [21]. Whereas the TiN and Nb(C, N) particles are kept relatively steady in the microstructure (Fig. 2d). Fig. 3 shows the typical engineering stress-strain curves of samples processed by hot rolling and solution treating. After solution treating at 1100 °C for 15 min, the ultimate tensile strength (UTS) of the sample is as high as 690 ± 11 MPa and the break elongation (EL) is 20 ± 2% due
3.2. Microstructural evolution of aged samples Fig. 4 shows the microstructural evolution of hot rolled samples during isothermal aging. It is obvious that a great deal of low angle grain boundaries are observed and recrystallization process is not activated, because the completed recrystallization temperature of hot rolled SFFSs is N1050 °C [12,13]. After 800 °C aging, the deformed microstructure characterized by dynamic recovery are reserved together with a further static recovery, which is also confirmed by many grain boundaries with misorientation angle of 2–5° inside the grains, as shown in Fig. 5. Nevertheless, the static recovery kinetics during the isothermal
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Fig. 2. EBSD images (a and c) and BSE (back scattered electron) images (b and d) showing the microstructure and precipitates of (a and b) the solution annealed and (c and d) the hot rolled plate, respectively. The SEM-EDS of TiN and Nb(C,N) are shown in (e) and (f). The grain size distribution of panel c is shown in panel (g). In (a) and (c), grain boundaries with misorientation larger than 15° are colored as black lines and the white lines represent low-angle grain boundaries with misorientation of 2–15°.
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(a)
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200μm Fig. 4. EBSD images showing the microstructure of the hot rolled samples after 800 °C aging treatment for different times: (a) 10 min, (b) 2 h and (c) 4 h. Grain boundaries with misorientation larger than 15° are colored as black lines. The white lines represent low-angle grain boundaries with misorientation of 2–15°.
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Fig. 5. Misorientation between grains in samples after (a) hot rolling and aging treatment at 800 °C for (b) 10 min, (c) 2 h and (d) 4 h. The distribution of low angle grain boundaries (b15°) is inserted in each image.
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aging is sluggish. For example, the fraction of low angle grain boundaries (b5°) in the sample aged for 10 min is about 75.0%, whereas the values for 2 h and 4 h increase slightly to ~76.9% and ~78.3%, respectively (Fig. 5). The average grain size of samples aged for 10 min, 2 h and 4 h, is about 65 ± 3, 67 ± 3 and 67 ± 4 μm, respectively. It should be noted that the growth of grains is not observed with increasing aging time, which is different from the results of solution treated (ST) samples. The grain size of aged samples without hot rolling increases from 67 ± 3 (ST) to 110 ± 7 μm (aged for 4 h). This contributes to the inhibition effect of precipitates on grain boundaries [14]. The precipitation of intermetallics is discussed in Section 3.3. Fig. 6 shows the microtexture of these samples suffering from different thermomechanical treatment. After hot rolling with reduction rate of ~25%, typical α-fiber (i.e., 〈110〉//RD, rolling direction) and γ-fiber (i.e., 〈111〉//ND, normal direction) texture are developed [23] and the
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peak value appears near {111} 〈123〉 component. After aging at 800 °C, the γ-fiber texture is weakened, while the α-fiber texture is strengthened. With increasing aging time, grain orientation is gradually transformed from γ-fiber texture to α fiber texture and the peak value evolves near {115} 〈110〉 component. 3.3. Effect of hot rolling on precipitation evolution Fig. 7 shows the TEM images and SAD of typical intermetallics in the samples aged at 800 °C. The results show that three kinds of intermetallics, Laves, chi, and sigma phase, are identified by TEM, which is in accord with the results predicted by JMatPro, as shown in the TTP curves embedded in Fig. 1. It is obvious from the TTP curves that precipitation of intermetallics phases, Laves, chi, and sigma, occur in the temperature
Fig. 6. Constant φ2 = 45° ODF(oriented distribution function) images showing the microtexture of samples after different thermomechanical processing: (a) solution treatment, (b) hot rolling, (c) aging at 800 °C for 10 min, (d) aging at 800 °C for 2 h and (e) aging at 800 °C for 4 h.
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Fig. 7. TEM images and SAD pattern of (a and b) Laves, (c and d) chi and (e and f) sigma phases in the hot rolled samples aged at 800 °C for (a and c) 10 min and (e) 2 h.
range 600–1000 °C. The kinetics curves of these precipitation display a C-shape with the “nose” temperature located between 800 and 950 °C. The TTP curves also show that the nucleation of Laves phase is faster than chi or sigma phase, and the formation temperature of Laves and chi phases is higher than that of sigma phase. The “nose” temperatures located between 800 and 950 °C of the calculated TTP curves in this research are higher than that in a 25Cr-3Mo-4Ni steel [8]. This is because that the high content of Cr and Mo enhances the stability of chi, sigma and Laves phases [5–8,13,14]. Whereas, the “nose” temperatures of these phases in 25Cr-3Mo-4Ni steels are uniform, which differs from the data in this paper. This contributes to the difference of chemical composition between 25Cr-3Mo-4Ni and 27Cr-4Mo-2Ni steels in this study. The morphology and distribution of intermetallic phases as a function of aging time are illustrated in Fig. 8. In Figs. 7 and 8, the rectangular Laves phase nucleates originally at dislocations and dislocation walls (Fig. 7a), and then develops into needle-like particles with the size of 0.1–0.3 μm. They are dispersedly distributed through the microstructure with increasing aging time (Fig. 8e). In the early stage of isothermal aging, the precipitation of fine chi phase appears at grain boundaries (Figs. 7c and 8a). With an increase in aging time, the amount of chi phase particles increases constantly and they form continuous networks along grain boundaries, as shown in Fig. 8b and c. After aging for 4 h, chi phase particles are found both along shear bands and around the TiN particles (Fig. 8c, d and e). Shear bands are marked within yellow dotted ellipses in Fig. 8c. Bulk sigma phase particles were firstly found around chi precipitates along grain boundaries. After aging for 2 h, sigma phase particles occupy most grain boundaries (Fig. 8b) and expand
into the interior of grains with a dendritic morphology (marked by red line with an arrow in Fig. 8d). The TEM image in Fig. 7d also shows the dendritic morphology of sigma phase. As the aging time increases to 4 h, sigma phase precipitates are observed around chi phase particles along shear bands (marked by blue line with an arrow in Fig. 8c and d). In addition, a few bulk sigma particles are also found to be distributed in the vicinity of TiN particles and some isolated sigma phase particles are also observed in the interior of grains (Fig. 8c). It should be pointed out that a number of microcracks are also observed (Fig. 8d) due to the result of the extremely brittle nature of this phase [5,14]. The results of SEM-EDS analysis in Table 2 show that the elements of chi phase are same as that of sigma phase: Fe, Cr, Mo Si and Ni. However, chi phase contains more Mo and less Cr than that of sigma phase. Laves phase is composed of Fe, Cr, Mo, Nb, Si, and Ni, and is enriched in Nb compared with both sigma and chi phases. One major objective of this study is to clarify the effect of hot rolling on precipitation reactions that occur in a 27Cr-4Mo-2Ni SFSS. Fig. 9 displays the microstructure of solution treated samples after isothermal aging for comparing. The results show that chi phase only forms at grain boundaries (Fig. 9a and b). A few sigma phase particles are observed and randomly distributed inside the grains except for the main precipitation of sigma phase at grain boundaries (Fig. 9b). Laves phases are distributed in the interior of grains (Fig. 9a and b). It is well known that second phases in steels preferentially precipitate at crystal defects, such as the boundaries of grain, subgrain and particles of other phases, dislocations, and vacancies [22]. In the present study, chi phase precipitates in samples aged for 10 min appears at grain boundaries due to the low barrier for heterogeneous nucleation. With increasing aging time,
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Fig. 8. BSE images showing the morphology and distribution of typical intermetallic in the hot rolled samples after aging treatment at 800 °C for different times: (a) 10 min, (b) 2 h, (c, d and e) 4 h. The image (d) and (e) are the high magnification picture of the ellipse area in the image (c).
the diffusion of Cr and Mo atoms, which compose the main elements of chi phase is sufficient [17,24]. Thus, shear bands characterized by low angle boundaries provide extra sites for the heterogeneous nucleation of chi phase (Fig. 8e) due to the high energy in shear bands, which offer the necessary drive force for the precipitation of chi phase [25,26]. As a result, the chi phase precipitated at both grain boundaries and shear bands are detected during long time aging. Precipitation of
sigma phase preferentially appears around the chi phase particles at boundaries, because areas surrounding chi phase particles, such as grain boundaries and shear bands possess enough Mo and Cr atoms [17,27,28]. Stress concentration around TiN particles is produced during hot rolling processes because of the different deformation ability of the ferrite matrix and TiN particles. Thus, high interfacial energy is generated, which is in favour of the formation of chi phase and sigma phase.
Table 2 EDS results of typical precipitates in 27Cr-4Mo-2Ni SFSS aged at 800 °C (wt%). Elements
Fe
Cr
Mo
Nb
Si
Ni
Matrix Sigma phase Chi phase Laves phase
67.1 ± 1.1 61.3 ± 1.5 58.0 ± 0.8 53.6 ± 0.8
27.8 ± 1.3 31.8 ± 1.2 22.8 ± 0.9 20.2 ± 0.8
3.0 ± 0.4 4.9 ± 0.5 15.7 ± 1.0 10.6 ± 0.9
0.3 ± 0.1 – – 10.3 ± 0.8
0.1 ± 0.1 0.4 ± 0.3 1.4 ± 0.4 1.5 ± 0.4
1.7 ± 0.1 1.6 ± 0.4 2.1 ± 0.2 3.8 ± 1.1
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Fig. 9. BSE images showing the morphology and distribution of typical intermetallic in the solution treated samples after aging treatment at 800 °C for (a) 2 h and (b) 4 h.
Moreover, the area fraction of the sigma phase is calculated by Image Pro, as shown in Fig. 10. It is obvious that sigma phase precipitation is accelerated by hot-rolling deformation, especially at long time aging. The area fraction is increased from 5.3 ± 0.4% to 8.7 ± 0.7% when aged for 4 h at 800 °C. The dislocations induced by hot rolling are suitable nucleation sites for Laves precipitation, as shown in Fig. 7a. This phenomenon is also found in other high Cr ferritic steels [8,29]. The dislocations in the crystal provide a path for the diffusion of solution atoms, such as Nb in this study, which promotes the diffusion reactions [30–32]. In consequence, a large quantity of Laves phase is observed at dislocations and dislocation cells. The size of the Laves phase in the hot rolled samples after aging for 4 h is 0.1–0.3 μm. While the size of Laves phase in the samples without hot rolling is 0.5–1.5 μm, which is similar to that in other report [14]. This difference was because that the dislocations and substructure produced by hot-rolling deformation provided additional sites for the nucleation of Laves phase particles. 3.4. Effect of precipitation on microstructural evolution Many studies show that there is an interaction between microstructure and precipitation [33]. The second phases at the grain boundaries or subgrain boundaries impede the recovery and recrystallization of ferrite grains during annealing processes [33–35]. In the present study, recovery of the deformed microstructure is dominant during the 800 °C aging, as shown in Fig. 5. After aging for 10 min at 800 °C, formation of chi phase at grain boundaries and Laves phase at dislocation walls are observed (Figs. 7a and 8a). The migration of grain boundaries during the further long time aging is inhibited due to the pinning effect of fine chi phase particles [14,29]. It is obvious that the grain boundaries
are completely covered by the chi phase (Figs. 8 and 9). Besides, due to the presence of intermetallic phases at grain boundaries and the relative high aging temperature (800 °C), solute drag effects may provide an additional impact [29]. Thus, the average grain sizes of hot rolled samples after aging are nearly constant. In addition, the recovery kinetics is sluggish due to the inhibition of fine Laves phase precipitates on the movement of dislocations, as shown in Fig. 11a and b. The Laves phase particles are marked by a yellow arrow in Fig. 11a. In general, the deformation in the microstructure of ferritic steels is heterogeneous due to different deformability between α-fiber and γfiber grains, and a few shear bands are usually observed in the γ-fiber grains [34]. In addition, strong recovery of the α-fiber grains are usually observed due to the kinetic and thermodynamic stability after recovery annealing [35]. In the present study, many fine chi phase particles are found along shear bands in γ-fiber grains (Figs. 4b and 8e). During medium temperature aging, such as 800 °C in this study, it is possible to further initiate recovery for α-fiber grains by way of slip and climb of dislocations despite of the hindrance of fine Laves phase particles on the movement of dislocations [14]. The subgrains are clearly observed in the TEM image of samples aged for 10 min, and the inner section of subgrain is relatively smooth, as shown in Fig. 11a. In contrast, the development of γ-fiber grains is weak due to the additional pinning effect of fine chi phase precipitates on shear bands although their storage energy is high [36,37]. The kernel average misorientation (KAM) in Fig. 11d also shows the higher energy in shear bands (higher KAM values). Therefore, recovery grains are rarely observed at shear bands (marked by black arrows), as shown in Fig. 11b and c. As a result, the α-fiber texture is strengthened slightly, while the γ-fiber texture is weakened with increasing aging time, as shown in Fig. 6. This finding agrees with the results in a FeCrAl stainless steel reported by Z. Sun [34].
Area fraction of sigma phase/%
10
3.5. Effect of microstructure and precipitation on microhardness
Hot rolled+Aged Solution treated+Aged
8 6 4 2 0
2h
4h Aging time/h
Fig. 10. Histogram showing the effect of hot rolling on area fraction of sigma phase after 800 °C aging.
For Vickers hardness testing, the average value of each specimen is calculated and the results are shown in Fig. 12. It is obvious that the microhardness of hot rolled samples is higher than that of solution treated ones due to the work hardening induce by hot rolling. The microhardness of aged samples decreases compared with the hot rolled one, and then increases continuously with increasing aging time. Both the reduction of work hardening induced by the additional aging recovery and the loss of solution strengthening caused by the precipitation of intermetallics account for the slight decrease of the hardness of samples aged for 10 min. The continuous increasing of hardness with increasing aging time is mainly due to the fast increase of both the size and content of sigma precipitates [12,14], as shown in Fig. 8 that the breadth of sigma phase particles extends to 15 μm, and in Fig. 10 that the area fraction of sigma phase particles increases to 8.7 ± 0.7%. Besides, the Vickers
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Fig. 11. TEM image (a), BC (band contrast) map (b), IPF (inverse pole figure) colored map (c) and KAM (kernel average misorientation) map (d) showing the recovery in samples aged at 800 °C for (a) 10 min, and (b, c and d) 2 h. The white lines in image (c) represent low-angle grain boundaries with misorientation of 2–15°.
hardness of hot rolled samples after aging is much higher than that of these solution treated ones. In general, the Vickers hardness of ferritic steels is related to the area fraction of precipitates [12,14]. Thus, the larger number of precipitates account for the increase in value of hardness for the hot rolled samples. The main reason is that the precipitation of intermetallics is evidently promoted by the hot-rolling deformation in the microstructure, as shown in Fig. 10. The slope of the curves of the sample aged for 4 h is larger than that for 2 h, indicating that the increase of Vickers hardness is accelerated by long aging. This is because that many additional chi and sigma phases form along shear bands in samples aged for 4 h (Fig. 10). Especially, the slope of the hardness curve between 10 min and 2 h in hot rolled samples is less than that in aged samples without hot rolling. Residual work hardening in the 400 Hot rolled +Aged Solution treated + Aged
Vickers hardness/HV
350 300
HR
250 ST
200 150 100
HR/ST
10 min
2h
4h
Thermomechanical treatment Fig. 12. Vickers hardness of samples with different thermomechanical treatments. (ST: solution treated, HR: hot rolled).
sample with 10 min aging after hot rolling makes the value of hardness rise to 295 ± 11 HV and then leads to a relatively gentle slope compared with the ones aged without hot rolling. 3.6. Effect of microstructure and precipitation on mechanical properties Fig. 13 shows the comparison of tensile properties for samples with different thermomechanical treatment. It is obvious that both the tensile strength and elongation (EL) of the samples decrease dramatically with increasing aging time. While the tensile strength of aged samples with hot rolling is lower than that of the solution treated ones, as shown in Fig. 13a. After aging for 10 min, the EL of hot rolled samples is about 20 ± 2%, while it decreases to 4 ± 0.8% at a long aging time for 4 h. The EL of aged samples without hot rolling are similar to the values of the hot rolled ones, as shown in Fig. 13b. In general, the strengthening mechanisms of ferritic steels mainly include solution strengthening, work hardening and precipitation strengthening [13,22]. Compared with hot rolled samples, the tensile strength decreases with a reduction of ~120 MPa due to the loss of work hardening induced by the process of recovery when the sample is aged at 800 °C for 10 min. However, the value of EL is improved to the grade of the samples treated by solution, because limited chi phase particles are observed in the grain boundaries. With an increase in aging time, the tensile strength of hot rolled samples aged for 2 h decreases to 583 ± 8 MPa and further to 529 ± 8 MPa after 4 h aging. This is because that precipitation and coarsening of intermetallics consume plenty of Cr, Mo and Si atoms, which lead to heavy loss of the solution strengthening [13,14]. In addition, the higher loss of the solution strengthening induced by more precipitation of intermetallics also accounts for the decrease of the tensile strength of hot rolled samples after aging compared to the solution treated ones. The precipitation of intermetallics at grain boundaries,
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(b) 800
Tensile strength/MPa
750
Solution treated + Aged Hot rolled + Aged
HR
700 650 600 ST 550
Solution treated + Aged Hot rolled + Aged
ST
20 15 HR 10 5
500 450
30 25
Elongation/%
(a)
0 ST/HR
10 min
2h Thermomechanical treatment
4h
ST/HR
2h 4h 10 min Thermomechanical treatment
Fig. 13. Comparison of tensile properties for samples with different thermomechanical treatments: (a) tensile strength and (b) elongation. (ST: solution treated, HT: hot rolled).
such as chi and sigma phases accounts for the rapid decline of EL. During tensile testing, the samples aged for 4 h suffer from a sudden fracture because most grain boundaries are covered by brittle chi and sigma phases. Besides, the microcracks throughout the sigma phase present an additional pernicious effect on the deterioration of EL, as shown in Fig. 8d. Accordingly, the fracture surface of the samples after tensile testing is observed by SEM, and the results are exhibited in Fig. 14. Many dimples are detected in the fracture surface of the samples aged for 10 min (Fig. 14a), which indicates a ductile fracture mode. In contrast, river pattern and cleavage facets are predominant on the surface of samples aged for 2–4 h (Fig. 14b and c). This presents a typical mode of brittle fracture [12,38]. It should be noted that precipitation and coarsening of grain boundary phases lead to a sharp decline of EL and make the
breakage pattern transformed from ductile to brittle fracture when the aging time increases from 10 min to 4 h. 4. Conclusions A 27Cr-4Mo-2Ni super ferritic stainless steel was hot rolled and subjected to isothermal aging at 800 °C for times ranging from between 10 min and 4 h. The evolution of microstructure and precipitation, and their effects on mechanical properties are investigated, and a few conclusions are drawn as follows: 1. After hot rolling, recovered ferrite grains with many low angle grain boundaries and shear bands, and embedded TiN and Nb(C,
Fig. 14. SEM images showing the impact fracture morphology of tested steels after aging treatment at 800 °C for different time: (a) 10 min, (b) 2 h and (c) 4 h.
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N) particles are observed in the microstructure. The tensile strength and elongation are 731 ± 10 MPa and 14 ± 3% respectively. 2. After aging at 800 °C, Laves phase nucleates at the location of dislocations and is distributed inside the grains. The nanosized chi phase particles start to precipitate at grain boundaries after short time aging. These chi phases form networks along grain boundaries and nucleate along shear bands with prolonged aging time. The precipitation of sigma phase appears around chi phase particles which locate at grain boundaries, shear bands and around TiN particles. Sigma phase is the dominant constituent and shows a dendritic morphology with increasing aging time. 3. Precipitation of intermetallics is accelerated by hot-rolling deformation. Substructure induced by hot rolling provides additional nucleation sites for intermetallics precipitation including chi and sigma phases. Plenty of dislocations in recovered grains present preferential locations for Laves phase precipitation. Recovery of hot rolled plate during aging is impeded by the formation of chi and Laves phases at dislocation and shear bands, respectively. 4. Vickers hardness increases with increasing aging time due to the higher content of the intermetallic phases formed after longer time aging. The value of hardness of hot rolled samples after aging is higher than that of aged samples without hot rolling. The tensile strength and elongation decrease dramatically with prolonged aging. Precipitation and coarsening of grain boundary phases account for the sharp decline of elongation and makes the breakage pattern transformed from ductile to brittle fracture when the aging time increases from 10 min to 4 h. Author contribution statement The first author (Hui-Hu Lu) presented the research idea and experimental design, and wrote this manuscript. The second and third authors (Hong-Kui Guo and Yi Luo) assisted to complete the hot rolling and annealing experiment. The fourth author (Zhen-Guang Liu) put forward some instructive suggestions for the writing of this paper and modified the language. The fifth author (Wen-Qi Li) helped to fulfill the work of SEM and EBSD analysis. The sixth author (Jian-Chun Li) helped to carry out rolling experiment and calculate the TTP curves. The TEM analysis was completed by the corresponding author (Wei Liang), and he was the faculty adviser of the first author. Acknowledgements This work was supported by Projects of International Cooperation in Shanxi (Grant No. 201603D421026). References [1] D. Janikowski, W. Henricks, Super-ferritic Stainless Steels - The Cost Effective Answer for Heat Transfer Tubing, REF: MP07-016, World Congress on Desalination and Water Reuse, Gran Canarias, Spain, 2007 (October 21–26). [2] P.A. Olubamb, J.H. Potgieter, L. Cornish, Corrosion behaviour of superferritic stainless steels cathodically modified with minor additions of ruthenium in sulphuric and hydrochloric acids, Mater. Des. 30 (2009) 1451–1457. [3] N. Pessall, J.I. Nurminen, Development of ferritic stainless steels for use in desalination plants, Corrosion 30 (11) (2013) 381–392. [4] I.A. Franson, Mechanical properties of high purity Fe-26 Cr-1 Mo ferritic stainless steel, Metall. Trans. 5 (1974) 2257–2264. [5] K. Premachandra, M.B. Cartie, R.H. Eric, Effect of stabilising elements on formation of σ phase in experimental ferritic stainless steels containing 39%Cr, Mater. Sci. Technol. 8 (2013) 437–442. [6] M.A. Streicher, Microstructures and some properties of Fe-28%Cr-4%Mo alloys, Corrosion 30 (4) (1974) 115–124. [7] T.J. Nichol, Microstructures and some properties of Fe-28%Cr-4%Mo alloys, Metall. Trans. A. 8 (1977) 229–237. [8] E.L. Brown, M.E. Burnett, P.T. Purtscher, G. Krauss, Intermetallic phase formation in 25Cr-3Mo-4Ni ferritic stainless steel, Metall. Trans. A. 14 (1983) 791–800.
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