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Influence of dynamic strain aging pre-treatment on the low-cycle fatigue behavior of modified 9Cr–1Mo steel
International Journal of Fatigue 47 (2013) 83–89
Contents lists available at SciVerse ScienceDirect
International Journal of Fatigue journal homepage: www.elsevier.com/locate/ijfatigue
Influence of dynamic strain aging pre-treatment on the low-cycle fatigue behavior of modified 9Cr–1Mo steel HongWei Zhou a,b,⇑, YiZhu He b, Hui Zhang b, YuWan Cen c a
School of Materials Science and Engineering, Jiangsu Key Lab of Advanced Metallic Materials, Southeast University, Nanjing 211189, Jiangshu, PR China School of Materials Science and Engineering, Anhui Key Lab of Materials Science and Processing, Anhui University of Technology, Maanshan 243002, Anhui, PR China c School of Mechanical Engineering, Anhui University of Technology, Maanshan 243002, Anhui, PR China b
a r t i c l e
i n f o
Article history: Received 29 March 2012 Received in revised form 3 July 2012 Accepted 26 July 2012 Available online 4 August 2012 Keywords: Modified 9Cr–1Mo steel Low-cycle fatigue Dynamic strain aging Fracture feature Dislocation cells
a b s t r a c t The influence of dynamic strain aging (DSA) pre-treatment on the low-cycle-fatigue (LCF) behavior of modified 9Cr–1Mo steel was investigated at 550 °C. The DSA pre-treatment reduces the fatigue life, which is reflected on the fracture surface as multiple crack initiation. The samples pre-treated by DSA have higher peak tensile stress and positive mean stress effects, which is responsible for the lifetime reduction. The DSA pre-treatment does not change cross-slip mechanisms during mechanical cycling, compared without DSA process, but results in accelerating the microstructure transformation from lath to cells with low dislocation densities, which reduces the number of cycles to failure. Ó 2012 Elsevier Ltd. All rights reserved.
1. Introduction Modified 9Cr–1Mo steel (ASME Grade P91/T91) has been widely used as a structural material in ultra-supercritical thermal power plants and petrochemical industries [1,2]. It is also a candidate material for use in the cladding of nuclear fuel, heat exchangers, reactor pressure vessels, and primary pipes for the next generation nuclear reactors [3,4]. The components of high-temperature systems are often subjected to repeated thermal stresses, which is due to temperature gradients, occurring on heating and cooling during start-ups and shut-downs or during temperature transients [5]. Hence, it is important to consider low-cycle fatigue (LCF) in the design of components. Dynamic strain aging (DSA) is the phenomenon of interactions between the diffusing atoms (C, N, and Cr) and mobile dislocations, which results in repeated pinning of dislocations and thus leads to enhanced work hardening. Therefore, the DSA process is an important strengthening mechanism in steels. The main effects of DSA on mechanical properties of materials during cyclic tests are an inverse dependence of the peak tensile stress on strain rate, an unusual increase in cyclic hardening, and a serrated flow known as the Portevin-Le Chatelier (PLC) effect [6,7].
⇑ Corresponding author at: School of Materials Science and Engineering, Jiangsu Key Lab of Advanced Metallic Materials, Southeast University, Nanjing 211189, Jiangshu, PR China. Tel.: +86 555 2311871; fax: +86 555 2311570. E-mail address: [email protected] (H. Zhou). 0142-1123/$ - see front matter Ó 2012 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.ijfatigue.2012.07.011
Some investigations have been reported about the influence of DSA on LCF behavior of austenitic stainless steels, and the results are contradictory [8–11]. The peak stress of the samples pre-treated by DSA was higher than that of solid-solution and cold working pre-treatment, but almost no differences of the fatigue life of the samples with and without the DSA treatment were found [9]. DSA in type 316L stainless steel has been found to enhance the stress response and reduce ductility, which localizes fatigue deformation, enhances fatigue cracking, and reduces fatigue life [8]. In contrast, in 2 1/4Cr–1Mo steels, an increase in the fatigue life was observed at the temperature range of 427–600 °C due to an interactive solid-solution hardening from the cyclic stress enhanced precipitation of Mo2C [10]. In 316L steel submitted to DSA pretreatment, the creep–fatigue cyclic strain amplitude is decreased and the material creep–fatigue life is effectively prolonged [11]. However, relative limited amount of research was reported about the effects of DSA on the LCF of matensitic steels. Mannan et al. have shown that P91 steel displayed certain evidences for the occurrence of DSA in the temperature range 500–600 °C under LCF [1,6]. It was found that other 9–12%Cr steels (10CrN and 10CrNW) present serrations or PLC characteristics of DSA in the hysteresis loops, but flow stress serrations have not been observed in P91 and P92 steels, at the temperature of 550 °C with the applied strain rate of 2 103 s1 [12]. PLC effect deduced from DSA has already been observed in 9–12%Cr steels on tensile curves with the flow stress serrations at temperatures between 227 °C
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and 450 °C and for intermediate strain rates [12–14]. It has been found that the regime of DSA under LCF loading is consistent with that in tensile loading [15]. Therefore, it is reasonable that the DSA effects happen at the same temperature from 227 °C to 450 °C under LCF, and the DSA phenomenon would disappear or weaken greatly at temperatures greater than 450 °C due to the loss of solutes responsible for locking of dislocations [14]. The present investigation aims at studying the effect of DSA pre-treatment on the fatigue of the P91 steel, and the microstructural changes and fracture feature during cyclic loading at 550 °C.
2. Experimental The chemical composition of the alloy used in the present study is shown in Table 1. The received P91 steel was normalized at 1080 °C and tempered at 760 °C. A martensitic microstructure with prior austenite grains (nearly 25 lm) is shown in Fig. 1. A group of tensile tests were conducted at different temperatures from room temperature (RT) to 550 °C with a constant strain rate of 3 104 s1. The stress-strain curves in Fig. 2 show that serrated flow appears with a temperature between 250 °C and 500 °C. Types A and D serrations, which are defined in Ref. [16], are observed in the stress–strain curves. Type A serration occurred at low temperatures of 250 °C and 350 °C, and Types D serrations at high temperatures of 450 °C. Types A serration with lower frequency first occurred at low strains, followed by Type A serration with high frequency at high strains at temperatures of 380 °C and 400 °C. It could be concluded from Fig. 2 that the DSA of P91 steel is most effective at 380 °C, which is accordance with previous studies [13,14]. Flow stress serrations in the samples deformed are possibly attributed to interaction between the diffusing atoms (C, N, and Cr) and mobile dislocations [8,14], which leads to dislocation multiplication and a higher dislocation density. This trend results in an increase in the hardening of the material. Predeformation at room temperature (RT) also leads to cold-work hardening, but the strengthening effect is inferior to the DSA pre-treatment at the same pre-strain [17]. The DSA process was conducted as follows. First, the as-received (AR) specimens were deformed to some pre-strain of 2.0% by a tensile test with a constant strain rate of 3 104 s1 at 380 °C, the as-called DSA pre-treatment. Pre-deformation process at RT is similar to the DSA process except the deformed temperature. In this paper, three groups of samples, the as-received (AR) condition, pre-deformed condition at RT, and the DSA pre-treatment, were prepared for fatigue experiments. The dog-bone specimens with a 28 mm gauge length and 7 mm gauge diameter were machined. The specimen surface was polished along the longitudinal direction with the emery paper down to #800 in order to remove surface defects, such as machining marks and scratches. All the tests were carried out in air under a fully-reversed, total axial strain control mode (Det/2 = ±0.2–1.0%, where Det is the total strain amplitude) at a constant strain rate of 8 10 s3 at 550 °C. A symmetrical triangular strain-time waveform was employed, using an IEHF-EM200k1-070-0A testing system (Shimaduz) controlled by an extensometer (Epsilon). The temperature variation along the gauge length of the specimen in a three-zone-furnace did not exceed ±3 °C. The fatigue life, Nf, is defined as the number of cycles corresponding to a 25% drop in the tensile stress at half-time.
Fig. 1. Optical micrograph of the as-received steel.
Fig. 2. Stress–strain diagram of tensile tests at different temperature.
JSM-6360LV Scanning electron microscopy (SEM) was applied to examine the fracture surface of the specimen. Transmission electron microscopy (TEM) examinations were conducted with a Philips Tecnai12 microscope, operating at 200 kV. It was used to investigate the evolution of laths and carbides under LCF conditions at 550 °C. Samples for TEM were obtained from thin slices cut at a distance of 2 mm away from the fracture surface. The slices were electrolytically polished using a double jet device with an electrolyte solution of 80 ml of methanol and 20 ml of perchloric acid.
3. Results 3.1. Fatigue life and cyclic stress strain curve The lifetimes for three groups of samples are shown in Fig. 3. It can be seen that the strengthening effects by cold-working at RT and the DSA pre-treatment at high temperature were detrimental to the LCF of P91 steel. Compared with pre-deformed condition at RT, the DSA pre-treatment leads to higher life reductions. Therefore, detailed studies were carried out to observe the influence of the DSA pre-treatment on the LCF. Continuous softening behaviors
Table 1 Chemical composition of the received P91. Element wt%
C 0.100
N 0.032
Cr 8.440
Mo 0.940
Mn 0.388
Si 0.370
Nb 0.083
V 0.190
Fe Balance
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Fig. 3. Fatigue life at 550 °C, Det/2 = ±0.2–1.0%. AR, RT-pre, and DSA-pre denote the as-received, pre-deform at RT, and DSA pre-treatment samples, respectively.
Fig. 4. Cyclic stress response at 550 °C, Det/2 = ±0.3–0.7%.
are observed through the whole life at all testing conditions, except the initial few cycles hardening at low strain amplitude of Det/ 2 = ±0.3%, as shown in Fig. 4. Cyclic softening has been also reported in 9–12% steels [5,6,18]. DSA pre-treatment has not changed the cyclic softening of the P91 steel, but the samples pretreated by DSA have a higher tensile stress response because of hardening induced by DSA. Fig. 5 shows that the mean stress of the AR samples is always negative and lower than that of DSA pre-treated samples, which changes from positive to negative after about 200 cycles. The samples submitted to DSA and RT pretreatment give higher elastic strain amplitude and lower plastic strain amplitude as compared to AR material at a given total strain before about 40 and 200 cycles, respectively. However, after these cycles until the fatigue life ends, the pretreated samples have higher plastic strain amplitude, compared to AR material with lower plastic strain amplitude, as shown in Fig. 6, where the plastic amplitudes and elastic amplitudes are acquired from the hysteresis loops per cycle at zero loads. At the same time, the DSA treated samples have higher plastic strain amplitude than that of the RT treated samples through the whole life. A lifetime prediction model depending directly from the amount of plastic strain per cycle has been published in Ref. [19]. Qualitatively, the lifetime reduction could result from this lifetime prediction model because of higher plastic strain with the DSA pre-treatment.
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Fig. 5. Peak stress and mean stress at 550 °C, Det/2 = ±0.5%.
Fig. 6. Evolution of plastic strain and elastic strain amplitudes at 550 °C, Det/ 2 = ±0.5%.
stead of the interior of the specimen, designated as stage I. Crack propagation designated as Stage II is transgranular under all LCF conditions, evidenced by striations on the fracture surface in Figs. 7c, 9b and 10c. There is no evidence of intergranular damage which could be induced by creep. The sudden fracture region of stage III exhibits ductile characteristic of P91 steel, as shown in Figs. 7d and 10d. It can be noticed that in Fig. 7b the crack initiation occurs at an inclusion near the surface. The inclusion-type nucleation can be understood as cyclic slip localization due to stress concentration at the inclusion, leading to either decohesion of the inclusion–matrix interface or cracking of the inclusion [20]. It can be seen only one crack propagates on the fracture surface of the AR specimens and leads to fracture in Figs. 7b and 8b, while several initiation sites are observed on the fracture surface of the specimens subjected to DSA pre-treatment in Figs. 9a and 10b. At the same time, multiple crack branching and formation of secondary cracks can be identified in Fig. 9b, which is attributed to hardening induced by DSA pre-treatment. Tyre pattern can be found in the high strain range of Det = ±0.5% in Fig. 10c, which is characteristic of LCF. It could be explained that crack propagation encountered hard particles such as precipitates, under the reciprocal action of cyclic loading, the crack jump around the hard particle and then move forward, leaving a indentation of row parallel [21]. 3.3. Microstructure evolution
3.2. Fracture features Figs. 7–10 show that fatigue cracks in all tested samples initiated at the slip bands connected to the specimen free surface in-
A TEM image of the general microstructure of tempered martensitic materials is given in Fig. 11. The microstructure is composed of martensitic laths with some precipitates along the grain
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Fig. 7. SEM micrographs of the fracture surface of the as-received specimen at 550 °C, Det/2 = ±0.3%, and the black dashed cycle shows crack initiation in (a) and crack initiation to occur at an inclusion in (b), (c) striations, and (d) dimple fracture.
Fig. 8. SEM micrographs of the fracture surface of the as-received specimen at 550 °C, Det/2 = ±0.5%, the black-dashed cycle shows crack initiation in (a) and (b).
Fig. 9. SEM micrographs of the fracture surface of the DSA pre-treated specimen at 550 °C, Det/2 = ±0.3%, the black-dashed cycles show crack multiple initiation in (a), (b) striations, The dark arrows show crack branching.
boundaries. According to previous studies [22,23], the coarser precipitates lying along the prior austenite grain boundaries, and lath boundaries are M23C6 carbides, while the finely distributed intragranular particles are Nb(C,N) and V(C,N). A TEM picture with high
dislocation density in the interior of laths is enlarged in the Fig. 11b. Fig. 12 shows that the initial lath structure in the as-received specimen is partly transformed into a cell structure, which is
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Fig. 10. SEM micrographs of the fracture surface of the DSA pre-treated specimen tested at 550 °C, Det/2 = ±0.5%, the black-dashed cycles show multiple initiation in (a) and (b), (c) The dark arrows show tyre pattern, and (d) transition from striations to dimple fracture.
Fig. 11. Microstructure of the as-received steel, black-line and dashed cycles in (a) show carbides, and dark arrows in (b) show high dislocation density.
Fig. 12. TEM micrographs in the as-received specimen at Det/2 = ±0.5%, the dark arrows show coarse carbides.
highly heterogeneous, during mechanical cycling at 550 °C. This trend is consistent with previous report [22]. In Fig. 13a, the specimen by DSA pre-treatment after LCF develops a more well-defined cell structure, which is more homogeneous. During fatigue deformation, the dislocations undergo a shuttling motion and they try
to get rearranged into a lower energy configuration such as cells and walls by cross-slip [18,24]. The lath structure in steels containing V remained intact for long periods at temperatures up to 600 °C [12]. This effect was attributed to the lath interfacial pinning by V(C, N) precipitates, which coarsen very slowly [25]. The softening
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Fig. 13. TEM micrographs in the fracture specimen by DSA pre-treatment at Det/2 = ±0.5%, (a) dashed cycles show new precipitates, (b) the black arrows show dislocation walls and the white arrows show little dislocations in the cells. Dashed cycles show coarse carbides.
behavior in martensitic stainless steels indicates that the martensitic lath structure of the steel is destabilized generally and the coarsening of the carbide particles accelerates under cyclic strain conditions, as shown in Figs. 12b and 13b. Consequently, such coarsening would be expected to promote the breakdown of the lath structure evolving to a cell structure. In addition, as shown in Fig. 13b, the dislocation density in the specimen by DSA pre-treatment is much lower than that in the normalized and tempered sample (Fig. 11b). It may be due to the greater annihilation rate of dislocations than the generation rate during cyclic plastic deformation. Fig. 13a reveals that very fine carbides are distributed into intragranular particles. Generally, new fine precipitates were considered to pin the lath, harden the matrix and prevent the gradual softening of the sample. However, the annihilation of dislocations and structure recovery in martensitic steels are dominant to softening. These precipitates (Fig. 13a) and the coarse carbides (Fig. 13b) are probable responsible for the formation of tyre pattern (Fig. 10) under cyclic strain conditions.
mogeneity of deformation, and rapid crack propagation due to the DSA-induced hardening [7,8,26]. In these LCF tests of the P91 steel, a cell structure is developed independent of the DSA pre-treatment. This trend indicates the DSA pre-treatment does not change the cross-slip mode of cyclic deformation in the martensitic steels, compared with the austenitic stainless steels. However, it results in accelerating the microstructure transformation from lath to cells and subgrains structure as shown in Fig. 12. In the case of the DSA pre-treatment, the samples during the mechanical cycling suffer the higher tensile stress response in Fig. 5, which not only leads to an increase in the amount of cross-slip, which should reflect in an increase in the number of crack initiation sites on the fracture surface (Figs. 9 and 10), but also accelerates crack propagation. The positive mean stress effects for fatigue tests are also a possible reason for the lifetime reduction after the DSA pre-treatment. The existence of a positive mean stress for compressive creep–fatigue tests was considered as a explanation for this lifetime reduction due to compressive holding periods [27].
4. Discussion
5. Conclusions
The previous results highlight the fact that cyclic peak stress decreases continuously till the end of fatigue life and dislocation cell structures are developed in the LCF tests of the P91 steel. Cyclic softening in Fig. 4 is attributed to the change from the original lath structure to cells or equiaxed subgrains, the decrease of the dislocation density, and the coarsening of precipitates. The lower energy cells structure with low dislocation densities is deleterious to fatigue. It has been reported that the variation of slip systems between {1 1 0}h1 1 1i and {1 1 2}h1 1 1i, depending on the Schmid factor, was responsible for the structure change from lath to dislocation cell structure [4]. DSA pre-treatment resulted in a high stress in Fig. 5 and a high plastic strain in Fig. 6, with a more well-defined cell structure in Fig. 13a. Other observations in the P91 steel [18,22] showed that the cells microstructure was much more homogeneous and pronounced when the applied strain was high, after creep–fatigue with the introduction of tension or compression hold. In other words, the high applied strain resulted in a high stress and a high plastic strain, which quickened dislocation slip and variation of slip systems between {1 1 0}h1 1 1i and {1 1 2}h1 1 1i. In the LCF tests of the austenitic stainless steels, the mechanism of cyclic plastic deformation changes from the cross-slip mode in the non-DSA regime to planar-slip mode in the DSA regime, which results in the transformation of dislocation structures from a cell structure to a planar one. The fatigue life is reduced by ways of multiple crack initiation, which comes from the DSA-induced inho-
Low-cycle-fatigue tests were conducted at the temperature of 550 °C for the P91 martensitic steel with and without the introduction of DSA pre-treatment at 380 °C where the DSA of P91 steel was most effective. The results of the present research are presented as follows: 1. The DSA pre-treatment does not change the cross-slip mode of cyclic deformation, however, resulting in accelerating the microstructure transformation from laths to cells and subgrains with low dislocation densities. 2. The DSA pre-treatment reduces the fatigue life by means of multiple crack initiation as a result of the increase in the amount of the stress-assisted cross-slip, and rapid crack propagation due to the DSA-induced high peak tensile stress. 3. A positive mean stress exists for the samples pre-treated by DSA under cyclic strain conditions, which is also considered as a reason for fatigue lifetime reduction after the DSA pre-treatment. 4. The carbides coarsening and new fine precipitates are possibly responsible for the formation of tyre pattern on the fracture surface.
Acknowledgements This work was financially supported by The National High Technology Research and Development Program of China (863 Program, Grant No. 2009AA044802), and Program for Innovative Research
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Team in Anhui University of technology. The authors thank Mr. Peter K. Liaw, University of Tennessee, for providing language help. References [1] Nagesha A, Valsan M, Kannan R, Bhanu Sankara Rao K, Mannan SL. Influence of temperature on the low cycle fatigue behaviour of a modified 9Cr–1Mo ferritic steel. Int J Fatigue 2002;24(12):1285–93. [2] Fournier B, Sauzay M, Caës C, Noblecourt M, Mottot M. Analysis of the hysteresis loops of a martensitic steel: part I: study of the influence of strain amplitude and temperature under pure fatigue loadings using an enhanced stress partitioning method. Mater Sci Eng A 2006;437(2):183–96. [3] Lee HY, Lee SH, Kim JB, Lee JH. Creep–fatigue damage for a structure with dissimilar metal welds of modified 9Cr–1Mo steel and 316L stainless steel. Int J Fatigue 2007;29(9–11):1868–79. [4] Kim DW, Kim SS. Contribution of microstructure and slip system to cyclic softening of 9 wt.%Cr steel. Int J Fatigue 2012;36(1):24–9. [5] Fournier B, Sauzay M, Caës C, Noblecourt M, Mottot M, Bougault A, et al. Creep–fatigue–oxidation interactions in a 9Cr–1Mo martensitic steel. Part I: Effect of tensile holding period on fatigue lifetime. Int J Fatigue 2008;30(4):649–62. [6] Mannan SL, Valsan M. High-temperature low cycle fatigue, creep–fatigue and thermomechanical fatigue of steels and their welds. Int J Mech Sci 2006;48(2):160–75. [7] Hong S-G, Lee S-B. The tensile and low-cycle fatigue behavior of cold worked 316L stainless steel: influence of dynamic strain aging. Int J Fatigue 2004;26(8):899–910. [8] Hong S-G, Lee K-O, Lee S-B. Dynamic strain aging effect on the fatigue resistance of type 316L stainless steel. Int J Fatigue 2005;27(10–12):1420–4. [9] Peng K, Chen W, Qian K. Effect of dynamic strain aging on low cycle fatigue behaviour of 18-8 austenitic stainless steel. Acta Metall Sinica Ser A 1993;6(6):420–5 [English edition]. [10] Challenger K, Miller A, Brinkman C. An explanation for the effects of hold periods on the elevated temperature fatigue behavior of 2 1/4 Cr–1 Mo steel. J Eng Mater Technol 1981;103(1):7–14. [11] Jiang H, Chen X, Fan Z, Dong J, Jiang H, Lu S. Influence of dynamic strain aging pre-treatment on creep–fatigue behavior in 316L stainless steel. Mater Sci Eng A 2009;500(1–2):98–103. [12] Fournier B, Dalle F, Sauzay M, Longour J, Salvi M, Caës C, et al. Comparison of various 9–12%Cr steels under fatigue and creep–fatigue loadings at high temperature. Mater Sci Eng A 2011;528(22–23):6934–45.
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Contents lists available at SciVerse ScienceDirect
International Journal of Fatigue journal homepage: www.elsevier.com/locate/ijfatigue
Influence of dynamic strain aging pre-treatment on the low-cycle fatigue behavior of modified 9Cr–1Mo steel HongWei Zhou a,b,⇑, YiZhu He b, Hui Zhang b, YuWan Cen c a
School of Materials Science and Engineering, Jiangsu Key Lab of Advanced Metallic Materials, Southeast University, Nanjing 211189, Jiangshu, PR China School of Materials Science and Engineering, Anhui Key Lab of Materials Science and Processing, Anhui University of Technology, Maanshan 243002, Anhui, PR China c School of Mechanical Engineering, Anhui University of Technology, Maanshan 243002, Anhui, PR China b
a r t i c l e
i n f o
Article history: Received 29 March 2012 Received in revised form 3 July 2012 Accepted 26 July 2012 Available online 4 August 2012 Keywords: Modified 9Cr–1Mo steel Low-cycle fatigue Dynamic strain aging Fracture feature Dislocation cells
a b s t r a c t The influence of dynamic strain aging (DSA) pre-treatment on the low-cycle-fatigue (LCF) behavior of modified 9Cr–1Mo steel was investigated at 550 °C. The DSA pre-treatment reduces the fatigue life, which is reflected on the fracture surface as multiple crack initiation. The samples pre-treated by DSA have higher peak tensile stress and positive mean stress effects, which is responsible for the lifetime reduction. The DSA pre-treatment does not change cross-slip mechanisms during mechanical cycling, compared without DSA process, but results in accelerating the microstructure transformation from lath to cells with low dislocation densities, which reduces the number of cycles to failure. Ó 2012 Elsevier Ltd. All rights reserved.
1. Introduction Modified 9Cr–1Mo steel (ASME Grade P91/T91) has been widely used as a structural material in ultra-supercritical thermal power plants and petrochemical industries [1,2]. It is also a candidate material for use in the cladding of nuclear fuel, heat exchangers, reactor pressure vessels, and primary pipes for the next generation nuclear reactors [3,4]. The components of high-temperature systems are often subjected to repeated thermal stresses, which is due to temperature gradients, occurring on heating and cooling during start-ups and shut-downs or during temperature transients [5]. Hence, it is important to consider low-cycle fatigue (LCF) in the design of components. Dynamic strain aging (DSA) is the phenomenon of interactions between the diffusing atoms (C, N, and Cr) and mobile dislocations, which results in repeated pinning of dislocations and thus leads to enhanced work hardening. Therefore, the DSA process is an important strengthening mechanism in steels. The main effects of DSA on mechanical properties of materials during cyclic tests are an inverse dependence of the peak tensile stress on strain rate, an unusual increase in cyclic hardening, and a serrated flow known as the Portevin-Le Chatelier (PLC) effect [6,7].
⇑ Corresponding author at: School of Materials Science and Engineering, Jiangsu Key Lab of Advanced Metallic Materials, Southeast University, Nanjing 211189, Jiangshu, PR China. Tel.: +86 555 2311871; fax: +86 555 2311570. E-mail address: [email protected] (H. Zhou). 0142-1123/$ - see front matter Ó 2012 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.ijfatigue.2012.07.011
Some investigations have been reported about the influence of DSA on LCF behavior of austenitic stainless steels, and the results are contradictory [8–11]. The peak stress of the samples pre-treated by DSA was higher than that of solid-solution and cold working pre-treatment, but almost no differences of the fatigue life of the samples with and without the DSA treatment were found [9]. DSA in type 316L stainless steel has been found to enhance the stress response and reduce ductility, which localizes fatigue deformation, enhances fatigue cracking, and reduces fatigue life [8]. In contrast, in 2 1/4Cr–1Mo steels, an increase in the fatigue life was observed at the temperature range of 427–600 °C due to an interactive solid-solution hardening from the cyclic stress enhanced precipitation of Mo2C [10]. In 316L steel submitted to DSA pretreatment, the creep–fatigue cyclic strain amplitude is decreased and the material creep–fatigue life is effectively prolonged [11]. However, relative limited amount of research was reported about the effects of DSA on the LCF of matensitic steels. Mannan et al. have shown that P91 steel displayed certain evidences for the occurrence of DSA in the temperature range 500–600 °C under LCF [1,6]. It was found that other 9–12%Cr steels (10CrN and 10CrNW) present serrations or PLC characteristics of DSA in the hysteresis loops, but flow stress serrations have not been observed in P91 and P92 steels, at the temperature of 550 °C with the applied strain rate of 2 103 s1 [12]. PLC effect deduced from DSA has already been observed in 9–12%Cr steels on tensile curves with the flow stress serrations at temperatures between 227 °C
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and 450 °C and for intermediate strain rates [12–14]. It has been found that the regime of DSA under LCF loading is consistent with that in tensile loading [15]. Therefore, it is reasonable that the DSA effects happen at the same temperature from 227 °C to 450 °C under LCF, and the DSA phenomenon would disappear or weaken greatly at temperatures greater than 450 °C due to the loss of solutes responsible for locking of dislocations [14]. The present investigation aims at studying the effect of DSA pre-treatment on the fatigue of the P91 steel, and the microstructural changes and fracture feature during cyclic loading at 550 °C.
2. Experimental The chemical composition of the alloy used in the present study is shown in Table 1. The received P91 steel was normalized at 1080 °C and tempered at 760 °C. A martensitic microstructure with prior austenite grains (nearly 25 lm) is shown in Fig. 1. A group of tensile tests were conducted at different temperatures from room temperature (RT) to 550 °C with a constant strain rate of 3 104 s1. The stress-strain curves in Fig. 2 show that serrated flow appears with a temperature between 250 °C and 500 °C. Types A and D serrations, which are defined in Ref. [16], are observed in the stress–strain curves. Type A serration occurred at low temperatures of 250 °C and 350 °C, and Types D serrations at high temperatures of 450 °C. Types A serration with lower frequency first occurred at low strains, followed by Type A serration with high frequency at high strains at temperatures of 380 °C and 400 °C. It could be concluded from Fig. 2 that the DSA of P91 steel is most effective at 380 °C, which is accordance with previous studies [13,14]. Flow stress serrations in the samples deformed are possibly attributed to interaction between the diffusing atoms (C, N, and Cr) and mobile dislocations [8,14], which leads to dislocation multiplication and a higher dislocation density. This trend results in an increase in the hardening of the material. Predeformation at room temperature (RT) also leads to cold-work hardening, but the strengthening effect is inferior to the DSA pre-treatment at the same pre-strain [17]. The DSA process was conducted as follows. First, the as-received (AR) specimens were deformed to some pre-strain of 2.0% by a tensile test with a constant strain rate of 3 104 s1 at 380 °C, the as-called DSA pre-treatment. Pre-deformation process at RT is similar to the DSA process except the deformed temperature. In this paper, three groups of samples, the as-received (AR) condition, pre-deformed condition at RT, and the DSA pre-treatment, were prepared for fatigue experiments. The dog-bone specimens with a 28 mm gauge length and 7 mm gauge diameter were machined. The specimen surface was polished along the longitudinal direction with the emery paper down to #800 in order to remove surface defects, such as machining marks and scratches. All the tests were carried out in air under a fully-reversed, total axial strain control mode (Det/2 = ±0.2–1.0%, where Det is the total strain amplitude) at a constant strain rate of 8 10 s3 at 550 °C. A symmetrical triangular strain-time waveform was employed, using an IEHF-EM200k1-070-0A testing system (Shimaduz) controlled by an extensometer (Epsilon). The temperature variation along the gauge length of the specimen in a three-zone-furnace did not exceed ±3 °C. The fatigue life, Nf, is defined as the number of cycles corresponding to a 25% drop in the tensile stress at half-time.
Fig. 1. Optical micrograph of the as-received steel.
Fig. 2. Stress–strain diagram of tensile tests at different temperature.
JSM-6360LV Scanning electron microscopy (SEM) was applied to examine the fracture surface of the specimen. Transmission electron microscopy (TEM) examinations were conducted with a Philips Tecnai12 microscope, operating at 200 kV. It was used to investigate the evolution of laths and carbides under LCF conditions at 550 °C. Samples for TEM were obtained from thin slices cut at a distance of 2 mm away from the fracture surface. The slices were electrolytically polished using a double jet device with an electrolyte solution of 80 ml of methanol and 20 ml of perchloric acid.
3. Results 3.1. Fatigue life and cyclic stress strain curve The lifetimes for three groups of samples are shown in Fig. 3. It can be seen that the strengthening effects by cold-working at RT and the DSA pre-treatment at high temperature were detrimental to the LCF of P91 steel. Compared with pre-deformed condition at RT, the DSA pre-treatment leads to higher life reductions. Therefore, detailed studies were carried out to observe the influence of the DSA pre-treatment on the LCF. Continuous softening behaviors
Table 1 Chemical composition of the received P91. Element wt%
C 0.100
N 0.032
Cr 8.440
Mo 0.940
Mn 0.388
Si 0.370
Nb 0.083
V 0.190
Fe Balance
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Fig. 3. Fatigue life at 550 °C, Det/2 = ±0.2–1.0%. AR, RT-pre, and DSA-pre denote the as-received, pre-deform at RT, and DSA pre-treatment samples, respectively.
Fig. 4. Cyclic stress response at 550 °C, Det/2 = ±0.3–0.7%.
are observed through the whole life at all testing conditions, except the initial few cycles hardening at low strain amplitude of Det/ 2 = ±0.3%, as shown in Fig. 4. Cyclic softening has been also reported in 9–12% steels [5,6,18]. DSA pre-treatment has not changed the cyclic softening of the P91 steel, but the samples pretreated by DSA have a higher tensile stress response because of hardening induced by DSA. Fig. 5 shows that the mean stress of the AR samples is always negative and lower than that of DSA pre-treated samples, which changes from positive to negative after about 200 cycles. The samples submitted to DSA and RT pretreatment give higher elastic strain amplitude and lower plastic strain amplitude as compared to AR material at a given total strain before about 40 and 200 cycles, respectively. However, after these cycles until the fatigue life ends, the pretreated samples have higher plastic strain amplitude, compared to AR material with lower plastic strain amplitude, as shown in Fig. 6, where the plastic amplitudes and elastic amplitudes are acquired from the hysteresis loops per cycle at zero loads. At the same time, the DSA treated samples have higher plastic strain amplitude than that of the RT treated samples through the whole life. A lifetime prediction model depending directly from the amount of plastic strain per cycle has been published in Ref. [19]. Qualitatively, the lifetime reduction could result from this lifetime prediction model because of higher plastic strain with the DSA pre-treatment.
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Fig. 5. Peak stress and mean stress at 550 °C, Det/2 = ±0.5%.
Fig. 6. Evolution of plastic strain and elastic strain amplitudes at 550 °C, Det/ 2 = ±0.5%.
stead of the interior of the specimen, designated as stage I. Crack propagation designated as Stage II is transgranular under all LCF conditions, evidenced by striations on the fracture surface in Figs. 7c, 9b and 10c. There is no evidence of intergranular damage which could be induced by creep. The sudden fracture region of stage III exhibits ductile characteristic of P91 steel, as shown in Figs. 7d and 10d. It can be noticed that in Fig. 7b the crack initiation occurs at an inclusion near the surface. The inclusion-type nucleation can be understood as cyclic slip localization due to stress concentration at the inclusion, leading to either decohesion of the inclusion–matrix interface or cracking of the inclusion [20]. It can be seen only one crack propagates on the fracture surface of the AR specimens and leads to fracture in Figs. 7b and 8b, while several initiation sites are observed on the fracture surface of the specimens subjected to DSA pre-treatment in Figs. 9a and 10b. At the same time, multiple crack branching and formation of secondary cracks can be identified in Fig. 9b, which is attributed to hardening induced by DSA pre-treatment. Tyre pattern can be found in the high strain range of Det = ±0.5% in Fig. 10c, which is characteristic of LCF. It could be explained that crack propagation encountered hard particles such as precipitates, under the reciprocal action of cyclic loading, the crack jump around the hard particle and then move forward, leaving a indentation of row parallel [21]. 3.3. Microstructure evolution
3.2. Fracture features Figs. 7–10 show that fatigue cracks in all tested samples initiated at the slip bands connected to the specimen free surface in-
A TEM image of the general microstructure of tempered martensitic materials is given in Fig. 11. The microstructure is composed of martensitic laths with some precipitates along the grain
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Fig. 7. SEM micrographs of the fracture surface of the as-received specimen at 550 °C, Det/2 = ±0.3%, and the black dashed cycle shows crack initiation in (a) and crack initiation to occur at an inclusion in (b), (c) striations, and (d) dimple fracture.
Fig. 8. SEM micrographs of the fracture surface of the as-received specimen at 550 °C, Det/2 = ±0.5%, the black-dashed cycle shows crack initiation in (a) and (b).
Fig. 9. SEM micrographs of the fracture surface of the DSA pre-treated specimen at 550 °C, Det/2 = ±0.3%, the black-dashed cycles show crack multiple initiation in (a), (b) striations, The dark arrows show crack branching.
boundaries. According to previous studies [22,23], the coarser precipitates lying along the prior austenite grain boundaries, and lath boundaries are M23C6 carbides, while the finely distributed intragranular particles are Nb(C,N) and V(C,N). A TEM picture with high
dislocation density in the interior of laths is enlarged in the Fig. 11b. Fig. 12 shows that the initial lath structure in the as-received specimen is partly transformed into a cell structure, which is
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Fig. 10. SEM micrographs of the fracture surface of the DSA pre-treated specimen tested at 550 °C, Det/2 = ±0.5%, the black-dashed cycles show multiple initiation in (a) and (b), (c) The dark arrows show tyre pattern, and (d) transition from striations to dimple fracture.
Fig. 11. Microstructure of the as-received steel, black-line and dashed cycles in (a) show carbides, and dark arrows in (b) show high dislocation density.
Fig. 12. TEM micrographs in the as-received specimen at Det/2 = ±0.5%, the dark arrows show coarse carbides.
highly heterogeneous, during mechanical cycling at 550 °C. This trend is consistent with previous report [22]. In Fig. 13a, the specimen by DSA pre-treatment after LCF develops a more well-defined cell structure, which is more homogeneous. During fatigue deformation, the dislocations undergo a shuttling motion and they try
to get rearranged into a lower energy configuration such as cells and walls by cross-slip [18,24]. The lath structure in steels containing V remained intact for long periods at temperatures up to 600 °C [12]. This effect was attributed to the lath interfacial pinning by V(C, N) precipitates, which coarsen very slowly [25]. The softening
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Fig. 13. TEM micrographs in the fracture specimen by DSA pre-treatment at Det/2 = ±0.5%, (a) dashed cycles show new precipitates, (b) the black arrows show dislocation walls and the white arrows show little dislocations in the cells. Dashed cycles show coarse carbides.
behavior in martensitic stainless steels indicates that the martensitic lath structure of the steel is destabilized generally and the coarsening of the carbide particles accelerates under cyclic strain conditions, as shown in Figs. 12b and 13b. Consequently, such coarsening would be expected to promote the breakdown of the lath structure evolving to a cell structure. In addition, as shown in Fig. 13b, the dislocation density in the specimen by DSA pre-treatment is much lower than that in the normalized and tempered sample (Fig. 11b). It may be due to the greater annihilation rate of dislocations than the generation rate during cyclic plastic deformation. Fig. 13a reveals that very fine carbides are distributed into intragranular particles. Generally, new fine precipitates were considered to pin the lath, harden the matrix and prevent the gradual softening of the sample. However, the annihilation of dislocations and structure recovery in martensitic steels are dominant to softening. These precipitates (Fig. 13a) and the coarse carbides (Fig. 13b) are probable responsible for the formation of tyre pattern (Fig. 10) under cyclic strain conditions.
mogeneity of deformation, and rapid crack propagation due to the DSA-induced hardening [7,8,26]. In these LCF tests of the P91 steel, a cell structure is developed independent of the DSA pre-treatment. This trend indicates the DSA pre-treatment does not change the cross-slip mode of cyclic deformation in the martensitic steels, compared with the austenitic stainless steels. However, it results in accelerating the microstructure transformation from lath to cells and subgrains structure as shown in Fig. 12. In the case of the DSA pre-treatment, the samples during the mechanical cycling suffer the higher tensile stress response in Fig. 5, which not only leads to an increase in the amount of cross-slip, which should reflect in an increase in the number of crack initiation sites on the fracture surface (Figs. 9 and 10), but also accelerates crack propagation. The positive mean stress effects for fatigue tests are also a possible reason for the lifetime reduction after the DSA pre-treatment. The existence of a positive mean stress for compressive creep–fatigue tests was considered as a explanation for this lifetime reduction due to compressive holding periods [27].
4. Discussion
5. Conclusions
The previous results highlight the fact that cyclic peak stress decreases continuously till the end of fatigue life and dislocation cell structures are developed in the LCF tests of the P91 steel. Cyclic softening in Fig. 4 is attributed to the change from the original lath structure to cells or equiaxed subgrains, the decrease of the dislocation density, and the coarsening of precipitates. The lower energy cells structure with low dislocation densities is deleterious to fatigue. It has been reported that the variation of slip systems between {1 1 0}h1 1 1i and {1 1 2}h1 1 1i, depending on the Schmid factor, was responsible for the structure change from lath to dislocation cell structure [4]. DSA pre-treatment resulted in a high stress in Fig. 5 and a high plastic strain in Fig. 6, with a more well-defined cell structure in Fig. 13a. Other observations in the P91 steel [18,22] showed that the cells microstructure was much more homogeneous and pronounced when the applied strain was high, after creep–fatigue with the introduction of tension or compression hold. In other words, the high applied strain resulted in a high stress and a high plastic strain, which quickened dislocation slip and variation of slip systems between {1 1 0}h1 1 1i and {1 1 2}h1 1 1i. In the LCF tests of the austenitic stainless steels, the mechanism of cyclic plastic deformation changes from the cross-slip mode in the non-DSA regime to planar-slip mode in the DSA regime, which results in the transformation of dislocation structures from a cell structure to a planar one. The fatigue life is reduced by ways of multiple crack initiation, which comes from the DSA-induced inho-
Low-cycle-fatigue tests were conducted at the temperature of 550 °C for the P91 martensitic steel with and without the introduction of DSA pre-treatment at 380 °C where the DSA of P91 steel was most effective. The results of the present research are presented as follows: 1. The DSA pre-treatment does not change the cross-slip mode of cyclic deformation, however, resulting in accelerating the microstructure transformation from laths to cells and subgrains with low dislocation densities. 2. The DSA pre-treatment reduces the fatigue life by means of multiple crack initiation as a result of the increase in the amount of the stress-assisted cross-slip, and rapid crack propagation due to the DSA-induced high peak tensile stress. 3. A positive mean stress exists for the samples pre-treated by DSA under cyclic strain conditions, which is also considered as a reason for fatigue lifetime reduction after the DSA pre-treatment. 4. The carbides coarsening and new fine precipitates are possibly responsible for the formation of tyre pattern on the fracture surface.
Acknowledgements This work was financially supported by The National High Technology Research and Development Program of China (863 Program, Grant No. 2009AA044802), and Program for Innovative Research
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