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Effect of strain rate on high-temperature low-cycle fatigue of 17-4 PH stainless steels
Materials Science and Engineering A 390 (2005) 291–298
Effect of strain rate on high-temperature low-cycle fatigue of 17-4 PH stainless steels Jui-Hung Wu, Chih-Kuang Lin∗ Department of Mechanical Engineering, National Central University, Chung-Li 32054, Taiwan Received 5 August 2004; received in revised form 5 August 2004; accepted 20 August 2004
Abstract The effect of strain rate (10−2 , 10−3 and 10−4 s−1 ) on the low-cycle fatigue (LCF) behavior was investigated for 17-4 PH stainless steels in three different conditions at temperatures of 300–500 ◦ C. The cyclic stress response (CSR) for Condition A tested at 300 and 400 ◦ C showed cyclic hardening due to an influence of dynamic strain aging (DSA). An in situ precipitation-hardening effect was found to be partially responsible for the cyclic hardening in Condition A at 400 ◦ C. For H900 and H1150 conditions tested at 300 and 400 ◦ C, the CSR exhibited a stable stress level before a fast drop in load indicating no cyclic hardening or softening. At 500 ◦ C, cyclic softening was observed for all given material conditions because of a thermal dislocation recovery mechanism. The cyclic softening behavior in Conditions A and H900 tested at 500 ◦ C is attributed partially to coarsening of the Cu-rich precipitates. The LCF life for each material condition, tested at a given temperature, decreased with decreasing strain rate as a result of an enhanced DSA effect. At all given testing conditions, transgranular cracking was the common fatigue fracture mode. © 2004 Elsevier B.V. All rights reserved. Keywords: 17-4 PH stainless steel; Low-cycle fatigue; High temperatures; Dynamic strain aging
1. Introduction Precipitation-hardening stainless steels have been widely used as structural components in applications such as nuclear, chemical, aircraft, and naval industries due to their excellent mechanical properties, fabrication characteristics and corrosion resistance. Of the former, 17-4 PH stainless steel is currently one of the most commonly used alloys [1]. In general, maximum strength and hardness can be obtained by aging at 450–510 ◦ C, which promotes the precipitation of coherent copper-rich clusters [1,2]. Aging at higher temperatures (above 540 ◦ C) would result in the precipitation of incoherent f.c.c. copper-rich phases, lower strength and hardness, and an enhancement in toughness [1,2]. Most of the previous studies on 17-4 PH stainless steels were focused on an analysis of microstructure, mechanical and fatigue properties at room-temperature [2–4]. Little or no work has been done on the high-temperature low-cycle ∗
Corresponding author. Tel.: +886 3 426 7340; fax: +886 3 425 4501. E-mail address: [email protected] (C.-K. Lin).
0921-5093/$ – see front matter © 2004 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2004.08.063
fatigue (LCF) behavior of this alloy. As in high-temperature applications, the components are often subjected to cyclic thermal stresses due to temperature gradients during start-up, shutdown and transients. Therefore, LCF property is essentially needed in the safe life design of components made from this alloy. As part of a series of studies on the high-temperature mechanical and fatigue properties of 17-4 PH stainless steels [5–7], the objective of the present study is to characterize the LCF behavior of such materials at high temperatures. In this study, systematic experiments were conducted to investigate the influence of strain rate (10−2 , 10−3 and 10−4 s−1 ) on LCF life, cyclic stress response (CSR), and dislocation substructure at 300–500 ◦ C for variously heat-treated 17-4 PH stainless steels.
2. Experimental procedures The commercially available 17-4 PH stainless steels used in the current study were supplied by the vendor in the form
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Table 1 Room-temperature mechanical properties of 17-4 PH stainless steels in different aged conditions Condition
Ultimate tensile strength (MPa)
Yield strength (MPa)
Elastic modulus (GPa)
Elongation (in 25 mm) (%)
Hardness (HRc)
V-notch impact toughness (J)
Condition A H900 H1150
1018 1414 966
992 1387 880
199 223 196
13.4 12.5 18.4
32.0 44.5 31.5
67 21 75
of hot-rolled, solution-annealed bars. The chemical composition of this alloy (wt.%) is 15.18 Cr, 4.47 Ni, 3.47 Cu, 0.65 Mn, 0.38 Si, 0.2 (Nb + Ti), 0.15 Mo, 0.03 S, 0.02 C, 0.016 P and Fe (balance). Three heat treatments were applied to the specimens, i.e., as-received “Condition A,” peak-aged “Condition H900,” and overaged “Condition H1150.” For Condition A, specimens were heated to 1038 ◦ C (1900 ◦ F), held 0.5 h at heat and cooled in air. For Conditions H900 and H1150, specimens were first heat-treated by solution annealing and then aged at 482 ◦ C (900 ◦ F) for 1 h and 621 ◦ C (1150 ◦ F) for 4 h, followed by cooling in air. The mechanical properties at room-temperature for each condition are listed in Table 1. LCF tests in ambient air at 300, 400, and 500 ◦ C were carried out under total axial strain-control on a commercial closed-loop servo-hydraulic test machine equipped with a SiC-heated furnace. A commercial direct-contact hightemperature extensometer with a 12 mm gage length was used to perform the strain-control. The axial smooth-surface specimens had a uniform cylindrical gage section of 6 mm in diameter and 18 mm in length. The LCF tests were performed using a symmetrical triangular waveform (strain ratio, R = −1) at a constant strain amplitude of 0.6% and at strain rates of 10−2 , 10−3 and 10−4 s−1 until failure or a 20% drop in load from the reference stress level taken at the tenth cycle. Each fatigue specimen was held in furnace for 15 min to reach thermal equilibrium prior to the start of fatigue testing. Scanning electron microscopy (SEM) was used for characterization of the fracture surface morphology, crack initiation, and propagation behavior in failed specimens. Microstructures of fatigue specimens were analyzed with transmission electron microscopy (TEM) to observe the dislocation substructure and precipitates. The thin foils for TEM analysis were prepared by cutting out 3 mm-diameter disks from some of the fatigue specimens followed by standard grinding and twin-jet electropolishing procedures.
3. Results and discussion 3.1. Cyclic stress response The CSR here was defined as the variation of the cyclic tensile peak stress with the number of cycles. Figs. 1–3 show the CSR curves under three strain rates (10−2 , 10−3 and 10−4 s−1 ) for variously heat-treated 17-4 PH stainless steels at 300, 400 and 500 ◦ C, respectively. It can be seen in Figs. 1 and 2 that at 300 and 400 ◦ C, the CSR for Condition
Fig. 1. Cyclic stress response curves at 300 ◦ C with various strain rates for 17-4 PH stainless steel in three conditions: (a) Condition A, (b) H900, and (c) H1150.
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400 ◦ C
Fig. 2. Cyclic stress response curves at with various strain rates for 17-4 PH stainless steel in three conditions: (a) Condition A, (b) H900, and (c) H1150.
A, under the given three strain rates, exhibited pronounced cyclic hardening behavior and a negative strain rate dependence of stress response, i.e., an increase in stress with decreasing strain rate. For H900 and H1150 conditions tested at 300 and 400 ◦ C, the CSR before a fast load drop did not show significant changes in peak stress level as observed for Condition A. The rapid drop in load on each CSR curve was believed to be caused by formation of microcracks and sub-
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Fig. 3. Cyclic stress response curves at 500 ◦ C with various strain rates for 17-4 PH stainless steel in three conditions: (a) Condition A, (b) H900, and (c) H1150.
sequent growth. However, at 500 ◦ C, the CSR curve for each material condition, under a given strain rate, exhibited gradually decreasing stress level (Fig. 3) indicative of cyclic softening behavior. The cyclic hardening behavior results from either an individual or combined effects of the following mechanisms: (i) dislocation generation and their interactions [8], (ii) formation of fine precipitates on dislocations during cyclic deformation [8], and (iii) interaction between mobile dislocations
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Table 2 Low-cycle fatigue results of Condition A at 300o C and at a strain amplitude of 0.6% Strain rate (s−1 ) 10−2 10−3 10−4
Half-life plastic strain range (%)
0.15 0.10 0.065
Tensile peak stress (MPa) First cycle (1)
Half-life (2)
Highest (3)
835 846 855
936 986 1053
961 1015 1065
and solute atoms, i.e., dynamic strain aging (DSA) [9]. In the present study, TEM analyses revealed no evidence of in situ precipitates in Condition A fatigue specimens tested at 300 ◦ C. A previous study [10] showed that solution-annealed Condition A needed a certain long period of time to complete the precipitation process when age-treated or tested at temperatures below 400 ◦ C. Our earlier work [5,6] indicated that the high-temperature yield strength of Condition A tested at 200 and 300 ◦ C remained unchanged as the testing time varied from 0.25 to 32 h. Apparently, the duration of a LCF test in the current work was not long enough for precipitation on dislocations to occur for Condition A tested at 300 ◦ C. Therefore, the role of in situ precipitation is excluded for causing cyclic hardening for Condition A at 300 ◦ C. On the other hand, LCF results of Condition A tested at 300 ◦ C revealed evidence for the occurrence of DSA. The CSR and the ratio of highest tensile peak stress to tensile peak stress of the first cycle both increased with a decrease in strain rate from10−2 to 10−4 s−1 , as shown in Fig. 1 and Table 2. These two characteristics are regarded as typical evidences of a DSA process in LCF [11]. In addition, it is seen in Table 2 that there was a continuous reduction in plastic strain range of the half-life cycle with a decrease in strain rate for Condition A at 300 ◦ C. This serves as an additional indication for the presence of DSA in LCF [11]. Hence, cyclic hardening for Condition A tested at 300 ◦ C is attributed to an interaction between mobile dislocations and diffusing solute atoms, i.e., DSA effect, during cyclic deformation. In the DSA regime, due to the immediate locking of dislocations by solute atoms, more new dislocations would be generated to maintain the imposed strain rate. This would cause a higher dislocation density and a continuous increase in stress with cycling. However, the serrated flow (PortevinLe Chatelier effect), a well-known manifestation of DSA, in the stress–strain hysteresis loops for Conditions A and H900 performed at 300 and 400 ◦ C at the different strain rates was not observed. An earlier work by Mannan [9] indicated a critical strain for the onset of serrations in flow curves was required. Consequently, the applied strain amplitude was considered to be lower than the critical strain value for appearance of serrations such that the stress–strain hysteresis loops for Conditions A and H900 tested at 300 and 400 ◦ C were lacking in serrations. It is concluded from these results that DSA took place before the strain level reached the critical value to generate serrated flow in stress–strain hysteresis loops.
Peak stress ratio (3)/(1)
Cycles to failure
1.151 1.200 1.246
674 449 367
For Condition A tested at 400 ◦ C, in addition to the influence of DSA, an in situ precipitation was also expected to play a role in causing cyclic hardening. Our previous investigation [5] on the high-temperature mechanical properties and high-cycle fatigue behavior of this alloy indicated that when unaged Condition A tested at 400 ◦ C, Cu-rich phases would gradually precipitate in the matrix and make the yield strength and fatigue resistance superior to tests of the lower temperatures. This in situ precipitation of Cu-rich phases contributes to the cyclic hardening behavior of Condition A tested at 400 ◦ C. However, the fine, coherent copper-rich precipitates expected at this stage could not be detected by TEM since the atomic scattering amplitude for electrons and atomic size are almost the same for Fe and Cu. Previous studies [3,5,6,10] have similarly shown that the copper-rich phase formed during the initial stage of precipitation process was a fine and coherent structure and did not generate any obvious strain contrast around it such that it was difficult to be detected by TEM. Nevertheless, it is still believed that in situ precipitation of Cu-rich phases is partially responsible for the cyclic hardening of Condition A at 400 ◦ C. As for the age-treated conditions, H900 and H1150, tested at 300 and 400 ◦ C, the aforementioned, typical features of DSA in LCF, such as a negative strain rate sensitivity, a greater extent of cyclic hardening at a lower strain rate, and a reduction of plastic strain range with decreasing strain rate, were not clear or even nonexistent, in particular for the H1150 condition. However, serrated flow, which characterized DSA in LCF, was observed in the stress–strain hysteresis loops for H1150 tested at all given conditions, as exemplified in Fig. 4. Apparently, the occurrence of DSA did not seem to influence CSR for the H900 and H1150 conditions. In other words, the intensity of DSA effect in the aged conditions, H900 and H1150, was less than that in the unaged Condition A. For steels, the DSA generally resulted from the interaction between diffusing carbon atoms and dislocations, particularly in the lower temperature region (50–200 ◦ C) [12,13]. TEM with energy dispersion spectrum (EDS) analyses indicated NbC carbides were observed in the unaged condition and the size (about 0.1 m) and amount of carbides were similar to those in aged conditions, H900 and H1150, as exemplified in Fig. 5. That is to say, the subsequent aging treatments did not change the morphology of carbides initially existing in the unaged condition. Therefore, the content of carbon in the matrix was barely reduced in the aged H900 and H1150 conditions. Hence, it is inferred that the diffusing carbon atoms are not responsible for the DSA effect in this steel. On the other hand,
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Fig. 4. Serrated flow in stress–strain hysteresis loops for Condition H1150 tested at 400 ◦ C with a strain rate of 10−2 s−1 .
the major difference between the unaged Condition A and the aged conditions, H900 and H1150, was the precipitation of Cu-rich phases. Note that Conditions H900 and H1150 have been initially age-treated at 482 and 621 ◦ C, respectively, to precipitate Cu-rich phases in the matrix prior to fatigue testing. From these results it is suggested that DSA was likely to be associated with an interaction of the Cu solute atoms and mobile dislocations. For H900 and H1150 conditions, due to the precipitation of Cu-rich phases in the initial aging treatment, the content of Cu solute atoms in the matrix would become lower and the intensity of DSA effect would be reduced. Hence, CSR of these two conditions at 300 and 400 ◦ C generally exhibited a stable, flat curve (Figs. 1 and 2). In order to further understand the cause of DSA, monotonic tensile tests for 17-4 PH alloy in the three conditions were conducted with a stroke rate of 0.5 mm/min at 300–500 ◦ C. The representative stress–strain curves of the three conditions at 400 ◦ C are shown in Fig. 6. It is seen in Fig. 6 that at 400 ◦ C serrated flow is clearly present in Condition A. However, in the agetreated H900 and H1150 conditions, where Cu-rich phases have been precipitated in the matrix prior to testing, the size of serrations was much smaller. An earlier study [9] indicated the mechanism of serrated flow in the low temperature region (below about 350 ◦ C) generally results from a diffusion of interstitial solutes to dislocations, while substitutional solutes are responsible for serrated flow in high-temperature region (above about 350 ◦ C). For this reason, substitutional Cu solute atoms are proposed to give rise to the DSA in this alloy. At 500 ◦ C, 17-4 PH alloys in the given three conditions exhibited cyclic softening under each given strain rate, as shown in Fig. 3. Cyclic softening of steels has been found
Fig. 5. TEM micrographs of carbides in 17-4 PH stainless steels: (a) Condition A and (b) Condition H900.
to arise from: (i) an annihilation of dislocations with a net decrease in dislocation density, (ii) rearrangement of dislocations into the configuration of cells and subgrains, or (iii) degradation of strength due to coarsening of precipitates [14]. TEM observations indicated that at 300 and 400 ◦ C the majority of grains in each material condition developed welldefined planar slip bands and evidence for this is shown in Fig. 7. Mechanism of interaction between solute atoms and dislocations during DSA would restrict cross slip of dislocations and enhance the slip planarity [15]. At 500 ◦ C the substructure consisted of both planar slip bands and cells with relatively lower dislocation density in the cell interiors and a uniform dispersion of Cu-rich precipitates in the matrix, as shown in Fig. 8. Apparently, thermal activation of dislocations became more operative than the hardening DSA effect for the given alloys tested at 500 ◦ C resulting in cyclic soft-
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Fig. 6. Monotonic stress–strain curves of 17-4 PH stainless steels in three conditions tested at 400 ◦ C.
ening. Moreover, the authors’ previous investigations [6,7] indicated that coarsening of Cu-rich phases was observed in Conditions A and H900 tested at 500 ◦ C and resulted in a reduction of hardness and fatigue resistance. Consequently, in addition to dynamic dislocation recovery, the coarsening of Cu-rich precipitates is believed to contribute to cyclic softening in Conditions A and H900 at 500 ◦ C. Furthermore, due to an increase in the period of a cycle with decreasing strain rate, the coarsening effect would become more effective such that the stress level was significantly decreased with
Fig. 8. TEM micrographs of microstructure in an H900 fatigue specimen tested at 500 ◦ C with a strain rate of 10−4 s−1 : (a) cell structures and (b) Cu-rich precipitates.
a reduction in strain rate for Conditions A and H900 tested at 500 ◦ C. 3.2. Effect of strain rate on fatigue life
Fig. 7. Planar slip bands formed in a Condition A fatigue specimen tested at 300 ◦ C with a strain rate of 10−4 s−1 .
Fig. 9 shows the variation of fatigue life as a function of strain rate at different temperatures for the heat-treated 17-4 PH stainless steels. It is seen in Fig. 9 that fatigue life of each condition, tested at a given temperature, decreased with a reduction in strain rate from 10−2 to 10−4 s−1 as a result of an enhanced DSA effect. DSA is considered to promote nonuniform fatigue deformation leading to localization of plastic flow and formation of intense slip bands [9,15]. In this regard, it would accelerate both crack initiation and
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Fig. 9. Variation of fatigue life with strain rate at various temperatures for 17-4 PH in different conditions: (a) Condition A, (b) H900, and (c) H1150.
propagation resulting in a significant reduction in fatigue life [9,15]. Moreover, under the DSA influence, the higher stress response developed during cyclic deformation caused a larger concentration of stress at the crack tip and an increase in crack growth rate leading to shorter crack propagation stage [9,11]. Generally, in the situations where DSA is present, the effect of DSA would be pronounced at decreasing strain rates or increasing temperatures [9]. Therefore, a decrease in strain rate from 10−2 to 10−4 s−1 , at a given temperature, makes the
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DSA effect far more effective to promote the development of planar slip bands and reduce the resistance to crack initiation and propagation leading to decreasing fatigue life. It has also been recognized by Srinivasan et al. [15] that a reduction in LCF life for a 316(N) stainless steel at 500–600 ◦ C, with decreasing strain rate, can be ascribed to the accelerated crack initiation and propagation caused by DSA. In addition, at 500 ◦ C and at a strain rate of 10−4 s−1 , a marked reduction in fatigue life of Conditions A and H900 in comparison with fatigue life at strain rates of 10−2 and 10−3 s−1 can be attributed to a coarsening of Cu-rich phases, which generates a degradation in material strength and greater plastic strain in each cycle. Fig. 9 indicates that for Conditions A and H900 the fatigue life in cycles, at a given strain rate, is not generally decreased with increasing temperature. In particular, at strain rates of 10−2 and 10−3 s−1 , fatigue life in either material condition has the highest value at the highest given temperature, 500 ◦ C. However, for the H1150 condition, the fatigue life at a given strain rate was generally independent of testing temperature between 300 and 500 ◦ C. It is believed that thermal dislocation recovery would induce greater material ductility and increase fatigue life under strain-control testing in terms of ductility exhaustion [16,17]. Therefore, in the current study, as testing temperature increased from 300 to 500 ◦ C, the enhanced thermal activation of dislocations is expected to increase material ductility and the stress needed to attain the imposed strain range decreases with increasing temperature. Hence, the fatigue lives at the higher temperatures can be effectively extended to a level close to or greater than at the lower temperature. An earlier study by Sandhya et al. [18] also indicated that LCF life of Alloy D9 increased with increasing temperature from 25 to 400 ◦ C due to reduction of stress needed to attain the applied strain range. However, the shorter fatigue lives of Conditions A and H900 tested at 500 ◦ C with a strain rate of 10−4 s−1 , as compared with the corresponding ones at 300 and 400 ◦ C, is attributed to the degradation of material strength and larger inelastic strain in each cycle caused by a pronounced coarsening effect of Curich precipitates. The total testing time was much longer at 10−4 s−1 strain rate such that coarsening effect of Cu-rich precipitates is more effective compared at the higher strain rates. SEM analyses indicated that at all given testing conditions, the fatigue fracture modes were commonly transgranular cracking accompanied by ductile striations on the fracture surface as shown in Fig. 10. Another example of transgranular cracking mode is given in Fig. 11 for a surface crack on the circumferential surface of a Condition H1150 specimen tested at 500 ◦ C and at a strain rate of 10−3 s−1 . In addition, no creep damages, such as cavities or intergranular microcracks, were observed for the given testing conditions. It has been reported that DSA would facilitate transgranular crack initiation and propagation in its operation regime [15]. Therefore, the crack initiation and propagation modes were transgranular at all of the given testing temperatures and stain rates.
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level followed by a rapid load drop due to formation of microcracks and their subsequent growth. 3. 17-4 PH alloys in the given three conditions at 500o C exhibited cyclic softening during LCF testing as a result of thermal dislocation recovery activities. Moreover, the cyclic softening behavior for Conditions A and H900 tested at 500o C is ascribed to the coarsening of Cu-rich precipitates. 4. For variously heat-treated 17-4 PH stainless steels tested at a given temperature, the number of cycles to failure consistently decreased with a reduction in strain rate due to an enhanced DSA effect. 5. At all given testing conditions, the fatigue failure for the given three material conditions occurred in a transgranular fracture mode.
Acknowledgement This work was funded by the National Science Council of the Republic of China (Taiwan) under Contract No. NSC-912216-E-008-007.
References Fig. 10. SEM fractographs of Condition H900 failed at 300 ◦ C with a strain rate of 10−3 s−1 : (a) initiation region and (b) stable crack growth region.
Fig. 11. SEM micrograph of a surface crack in a Condition H1150 fatigue specimen tested at 500 ◦ C with a strain rate of 10−3 s−1 .
4. Conclusions 1. The marked cyclic hardening behavior for Condition A tested at 300 and 400o C is attributed to dynamic strain aging effect during fatigue deformation. In addition, the in situ precipitation of Cu-rich phases is partially responsible for cyclic hardening in Condition A at 400o C. 2. The cyclic stress response of H900 and H1150 conditions tested at 300 and 400o C generally exhibited a stable stress
[1] W.F. Smith, Structure and Properties of Engineering Alloys, 2nd ed., McGraw-Hill, New York, 1993, pp. 328–335. [2] H.J. Rack, D. Kalish, Metall. Trans. 5 (1974) 1595. [3] U.K. Viswanathan, S. Banerjee, R. Krishnan, Mater. Sci. Eng. A 104 (1988) 181. [4] U.K. Viswanathan, P.K.K. Nayar, R. Krishnan, Mater. Sci. Tech. 5 (1989) 346. [5] J.-H. Wu, C.-K. Lin, Metall. Mater. Trans. A 33A (2002) 1715. [6] J.-H. Wu, C.-K. Lin, J. Mater. Sci. 38 (2003) 965. [7] J.-H. Wu, C.-K. Lin, Mater. Trans. 44 (2003) 713. [8] K.B.S. Rao, H. Schiffers, H. Schuster, Metall. Trans. A 19A (1988) 359. [9] S.L. Mannan, Bull. Mater. Sci. 16 (1993) 561. [10] U.K. Viswanathan, P.K.K. Nayar, R. Krishnan, Mater. Sci. Tech. 5 (1989) 346. [11] V.S. Srinivasan, R. Sandhya, K.B.S. Rao, S.L. Mannan, K.S. Raghavan, Int. J. Fatigue 13 (1991) 471. [12] S. Herenu, I. Alvarez-Armas, A.F. Armas, Scrip. Mater. 45 (2001) 739. [13] M.J. Roberts, W.S. Owen, Metall. Trans. 1 (1970) 3203. [14] A. Nagesha, M. Valsan, R. Kannan, K.B.S. Rao, S.L. Mannan, Int. J. Fatigue 24 (2002) 1285. [15] V.S. Srinivasan, M. Valsan, R. Sandhya, K.B.S. Rao, S.L. Mannan, D.H. Sastry, Int. J. Fatigue 21 (1999) 11. [16] K.B.S. Rao, M. Valsan, R. Sandhya, S.L. Mannan, P. Rodriguez, Metall. Trans. 24A (1993) 913. [17] A. Plumtree, L. Pawlus, Substructural Developments During Strain Cycling of Wavy Slip Materials in Low Cycle Fatigue, ASTM STP 942, American Society for Testing and Materials, Philadelphia, PA, 1988, pp. 81–97. [18] R. Sandhya, K.B.S. Rao, S.L. Mannan, Scrip. Mater. 41 (1999) 921.
Effect of strain rate on high-temperature low-cycle fatigue of 17-4 PH stainless steels Jui-Hung Wu, Chih-Kuang Lin∗ Department of Mechanical Engineering, National Central University, Chung-Li 32054, Taiwan Received 5 August 2004; received in revised form 5 August 2004; accepted 20 August 2004
Abstract The effect of strain rate (10−2 , 10−3 and 10−4 s−1 ) on the low-cycle fatigue (LCF) behavior was investigated for 17-4 PH stainless steels in three different conditions at temperatures of 300–500 ◦ C. The cyclic stress response (CSR) for Condition A tested at 300 and 400 ◦ C showed cyclic hardening due to an influence of dynamic strain aging (DSA). An in situ precipitation-hardening effect was found to be partially responsible for the cyclic hardening in Condition A at 400 ◦ C. For H900 and H1150 conditions tested at 300 and 400 ◦ C, the CSR exhibited a stable stress level before a fast drop in load indicating no cyclic hardening or softening. At 500 ◦ C, cyclic softening was observed for all given material conditions because of a thermal dislocation recovery mechanism. The cyclic softening behavior in Conditions A and H900 tested at 500 ◦ C is attributed partially to coarsening of the Cu-rich precipitates. The LCF life for each material condition, tested at a given temperature, decreased with decreasing strain rate as a result of an enhanced DSA effect. At all given testing conditions, transgranular cracking was the common fatigue fracture mode. © 2004 Elsevier B.V. All rights reserved. Keywords: 17-4 PH stainless steel; Low-cycle fatigue; High temperatures; Dynamic strain aging
1. Introduction Precipitation-hardening stainless steels have been widely used as structural components in applications such as nuclear, chemical, aircraft, and naval industries due to their excellent mechanical properties, fabrication characteristics and corrosion resistance. Of the former, 17-4 PH stainless steel is currently one of the most commonly used alloys [1]. In general, maximum strength and hardness can be obtained by aging at 450–510 ◦ C, which promotes the precipitation of coherent copper-rich clusters [1,2]. Aging at higher temperatures (above 540 ◦ C) would result in the precipitation of incoherent f.c.c. copper-rich phases, lower strength and hardness, and an enhancement in toughness [1,2]. Most of the previous studies on 17-4 PH stainless steels were focused on an analysis of microstructure, mechanical and fatigue properties at room-temperature [2–4]. Little or no work has been done on the high-temperature low-cycle ∗
Corresponding author. Tel.: +886 3 426 7340; fax: +886 3 425 4501. E-mail address: [email protected] (C.-K. Lin).
0921-5093/$ – see front matter © 2004 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2004.08.063
fatigue (LCF) behavior of this alloy. As in high-temperature applications, the components are often subjected to cyclic thermal stresses due to temperature gradients during start-up, shutdown and transients. Therefore, LCF property is essentially needed in the safe life design of components made from this alloy. As part of a series of studies on the high-temperature mechanical and fatigue properties of 17-4 PH stainless steels [5–7], the objective of the present study is to characterize the LCF behavior of such materials at high temperatures. In this study, systematic experiments were conducted to investigate the influence of strain rate (10−2 , 10−3 and 10−4 s−1 ) on LCF life, cyclic stress response (CSR), and dislocation substructure at 300–500 ◦ C for variously heat-treated 17-4 PH stainless steels.
2. Experimental procedures The commercially available 17-4 PH stainless steels used in the current study were supplied by the vendor in the form
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Table 1 Room-temperature mechanical properties of 17-4 PH stainless steels in different aged conditions Condition
Ultimate tensile strength (MPa)
Yield strength (MPa)
Elastic modulus (GPa)
Elongation (in 25 mm) (%)
Hardness (HRc)
V-notch impact toughness (J)
Condition A H900 H1150
1018 1414 966
992 1387 880
199 223 196
13.4 12.5 18.4
32.0 44.5 31.5
67 21 75
of hot-rolled, solution-annealed bars. The chemical composition of this alloy (wt.%) is 15.18 Cr, 4.47 Ni, 3.47 Cu, 0.65 Mn, 0.38 Si, 0.2 (Nb + Ti), 0.15 Mo, 0.03 S, 0.02 C, 0.016 P and Fe (balance). Three heat treatments were applied to the specimens, i.e., as-received “Condition A,” peak-aged “Condition H900,” and overaged “Condition H1150.” For Condition A, specimens were heated to 1038 ◦ C (1900 ◦ F), held 0.5 h at heat and cooled in air. For Conditions H900 and H1150, specimens were first heat-treated by solution annealing and then aged at 482 ◦ C (900 ◦ F) for 1 h and 621 ◦ C (1150 ◦ F) for 4 h, followed by cooling in air. The mechanical properties at room-temperature for each condition are listed in Table 1. LCF tests in ambient air at 300, 400, and 500 ◦ C were carried out under total axial strain-control on a commercial closed-loop servo-hydraulic test machine equipped with a SiC-heated furnace. A commercial direct-contact hightemperature extensometer with a 12 mm gage length was used to perform the strain-control. The axial smooth-surface specimens had a uniform cylindrical gage section of 6 mm in diameter and 18 mm in length. The LCF tests were performed using a symmetrical triangular waveform (strain ratio, R = −1) at a constant strain amplitude of 0.6% and at strain rates of 10−2 , 10−3 and 10−4 s−1 until failure or a 20% drop in load from the reference stress level taken at the tenth cycle. Each fatigue specimen was held in furnace for 15 min to reach thermal equilibrium prior to the start of fatigue testing. Scanning electron microscopy (SEM) was used for characterization of the fracture surface morphology, crack initiation, and propagation behavior in failed specimens. Microstructures of fatigue specimens were analyzed with transmission electron microscopy (TEM) to observe the dislocation substructure and precipitates. The thin foils for TEM analysis were prepared by cutting out 3 mm-diameter disks from some of the fatigue specimens followed by standard grinding and twin-jet electropolishing procedures.
3. Results and discussion 3.1. Cyclic stress response The CSR here was defined as the variation of the cyclic tensile peak stress with the number of cycles. Figs. 1–3 show the CSR curves under three strain rates (10−2 , 10−3 and 10−4 s−1 ) for variously heat-treated 17-4 PH stainless steels at 300, 400 and 500 ◦ C, respectively. It can be seen in Figs. 1 and 2 that at 300 and 400 ◦ C, the CSR for Condition
Fig. 1. Cyclic stress response curves at 300 ◦ C with various strain rates for 17-4 PH stainless steel in three conditions: (a) Condition A, (b) H900, and (c) H1150.
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400 ◦ C
Fig. 2. Cyclic stress response curves at with various strain rates for 17-4 PH stainless steel in three conditions: (a) Condition A, (b) H900, and (c) H1150.
A, under the given three strain rates, exhibited pronounced cyclic hardening behavior and a negative strain rate dependence of stress response, i.e., an increase in stress with decreasing strain rate. For H900 and H1150 conditions tested at 300 and 400 ◦ C, the CSR before a fast load drop did not show significant changes in peak stress level as observed for Condition A. The rapid drop in load on each CSR curve was believed to be caused by formation of microcracks and sub-
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Fig. 3. Cyclic stress response curves at 500 ◦ C with various strain rates for 17-4 PH stainless steel in three conditions: (a) Condition A, (b) H900, and (c) H1150.
sequent growth. However, at 500 ◦ C, the CSR curve for each material condition, under a given strain rate, exhibited gradually decreasing stress level (Fig. 3) indicative of cyclic softening behavior. The cyclic hardening behavior results from either an individual or combined effects of the following mechanisms: (i) dislocation generation and their interactions [8], (ii) formation of fine precipitates on dislocations during cyclic deformation [8], and (iii) interaction between mobile dislocations
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Table 2 Low-cycle fatigue results of Condition A at 300o C and at a strain amplitude of 0.6% Strain rate (s−1 ) 10−2 10−3 10−4
Half-life plastic strain range (%)
0.15 0.10 0.065
Tensile peak stress (MPa) First cycle (1)
Half-life (2)
Highest (3)
835 846 855
936 986 1053
961 1015 1065
and solute atoms, i.e., dynamic strain aging (DSA) [9]. In the present study, TEM analyses revealed no evidence of in situ precipitates in Condition A fatigue specimens tested at 300 ◦ C. A previous study [10] showed that solution-annealed Condition A needed a certain long period of time to complete the precipitation process when age-treated or tested at temperatures below 400 ◦ C. Our earlier work [5,6] indicated that the high-temperature yield strength of Condition A tested at 200 and 300 ◦ C remained unchanged as the testing time varied from 0.25 to 32 h. Apparently, the duration of a LCF test in the current work was not long enough for precipitation on dislocations to occur for Condition A tested at 300 ◦ C. Therefore, the role of in situ precipitation is excluded for causing cyclic hardening for Condition A at 300 ◦ C. On the other hand, LCF results of Condition A tested at 300 ◦ C revealed evidence for the occurrence of DSA. The CSR and the ratio of highest tensile peak stress to tensile peak stress of the first cycle both increased with a decrease in strain rate from10−2 to 10−4 s−1 , as shown in Fig. 1 and Table 2. These two characteristics are regarded as typical evidences of a DSA process in LCF [11]. In addition, it is seen in Table 2 that there was a continuous reduction in plastic strain range of the half-life cycle with a decrease in strain rate for Condition A at 300 ◦ C. This serves as an additional indication for the presence of DSA in LCF [11]. Hence, cyclic hardening for Condition A tested at 300 ◦ C is attributed to an interaction between mobile dislocations and diffusing solute atoms, i.e., DSA effect, during cyclic deformation. In the DSA regime, due to the immediate locking of dislocations by solute atoms, more new dislocations would be generated to maintain the imposed strain rate. This would cause a higher dislocation density and a continuous increase in stress with cycling. However, the serrated flow (PortevinLe Chatelier effect), a well-known manifestation of DSA, in the stress–strain hysteresis loops for Conditions A and H900 performed at 300 and 400 ◦ C at the different strain rates was not observed. An earlier work by Mannan [9] indicated a critical strain for the onset of serrations in flow curves was required. Consequently, the applied strain amplitude was considered to be lower than the critical strain value for appearance of serrations such that the stress–strain hysteresis loops for Conditions A and H900 tested at 300 and 400 ◦ C were lacking in serrations. It is concluded from these results that DSA took place before the strain level reached the critical value to generate serrated flow in stress–strain hysteresis loops.
Peak stress ratio (3)/(1)
Cycles to failure
1.151 1.200 1.246
674 449 367
For Condition A tested at 400 ◦ C, in addition to the influence of DSA, an in situ precipitation was also expected to play a role in causing cyclic hardening. Our previous investigation [5] on the high-temperature mechanical properties and high-cycle fatigue behavior of this alloy indicated that when unaged Condition A tested at 400 ◦ C, Cu-rich phases would gradually precipitate in the matrix and make the yield strength and fatigue resistance superior to tests of the lower temperatures. This in situ precipitation of Cu-rich phases contributes to the cyclic hardening behavior of Condition A tested at 400 ◦ C. However, the fine, coherent copper-rich precipitates expected at this stage could not be detected by TEM since the atomic scattering amplitude for electrons and atomic size are almost the same for Fe and Cu. Previous studies [3,5,6,10] have similarly shown that the copper-rich phase formed during the initial stage of precipitation process was a fine and coherent structure and did not generate any obvious strain contrast around it such that it was difficult to be detected by TEM. Nevertheless, it is still believed that in situ precipitation of Cu-rich phases is partially responsible for the cyclic hardening of Condition A at 400 ◦ C. As for the age-treated conditions, H900 and H1150, tested at 300 and 400 ◦ C, the aforementioned, typical features of DSA in LCF, such as a negative strain rate sensitivity, a greater extent of cyclic hardening at a lower strain rate, and a reduction of plastic strain range with decreasing strain rate, were not clear or even nonexistent, in particular for the H1150 condition. However, serrated flow, which characterized DSA in LCF, was observed in the stress–strain hysteresis loops for H1150 tested at all given conditions, as exemplified in Fig. 4. Apparently, the occurrence of DSA did not seem to influence CSR for the H900 and H1150 conditions. In other words, the intensity of DSA effect in the aged conditions, H900 and H1150, was less than that in the unaged Condition A. For steels, the DSA generally resulted from the interaction between diffusing carbon atoms and dislocations, particularly in the lower temperature region (50–200 ◦ C) [12,13]. TEM with energy dispersion spectrum (EDS) analyses indicated NbC carbides were observed in the unaged condition and the size (about 0.1 m) and amount of carbides were similar to those in aged conditions, H900 and H1150, as exemplified in Fig. 5. That is to say, the subsequent aging treatments did not change the morphology of carbides initially existing in the unaged condition. Therefore, the content of carbon in the matrix was barely reduced in the aged H900 and H1150 conditions. Hence, it is inferred that the diffusing carbon atoms are not responsible for the DSA effect in this steel. On the other hand,
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Fig. 4. Serrated flow in stress–strain hysteresis loops for Condition H1150 tested at 400 ◦ C with a strain rate of 10−2 s−1 .
the major difference between the unaged Condition A and the aged conditions, H900 and H1150, was the precipitation of Cu-rich phases. Note that Conditions H900 and H1150 have been initially age-treated at 482 and 621 ◦ C, respectively, to precipitate Cu-rich phases in the matrix prior to fatigue testing. From these results it is suggested that DSA was likely to be associated with an interaction of the Cu solute atoms and mobile dislocations. For H900 and H1150 conditions, due to the precipitation of Cu-rich phases in the initial aging treatment, the content of Cu solute atoms in the matrix would become lower and the intensity of DSA effect would be reduced. Hence, CSR of these two conditions at 300 and 400 ◦ C generally exhibited a stable, flat curve (Figs. 1 and 2). In order to further understand the cause of DSA, monotonic tensile tests for 17-4 PH alloy in the three conditions were conducted with a stroke rate of 0.5 mm/min at 300–500 ◦ C. The representative stress–strain curves of the three conditions at 400 ◦ C are shown in Fig. 6. It is seen in Fig. 6 that at 400 ◦ C serrated flow is clearly present in Condition A. However, in the agetreated H900 and H1150 conditions, where Cu-rich phases have been precipitated in the matrix prior to testing, the size of serrations was much smaller. An earlier study [9] indicated the mechanism of serrated flow in the low temperature region (below about 350 ◦ C) generally results from a diffusion of interstitial solutes to dislocations, while substitutional solutes are responsible for serrated flow in high-temperature region (above about 350 ◦ C). For this reason, substitutional Cu solute atoms are proposed to give rise to the DSA in this alloy. At 500 ◦ C, 17-4 PH alloys in the given three conditions exhibited cyclic softening under each given strain rate, as shown in Fig. 3. Cyclic softening of steels has been found
Fig. 5. TEM micrographs of carbides in 17-4 PH stainless steels: (a) Condition A and (b) Condition H900.
to arise from: (i) an annihilation of dislocations with a net decrease in dislocation density, (ii) rearrangement of dislocations into the configuration of cells and subgrains, or (iii) degradation of strength due to coarsening of precipitates [14]. TEM observations indicated that at 300 and 400 ◦ C the majority of grains in each material condition developed welldefined planar slip bands and evidence for this is shown in Fig. 7. Mechanism of interaction between solute atoms and dislocations during DSA would restrict cross slip of dislocations and enhance the slip planarity [15]. At 500 ◦ C the substructure consisted of both planar slip bands and cells with relatively lower dislocation density in the cell interiors and a uniform dispersion of Cu-rich precipitates in the matrix, as shown in Fig. 8. Apparently, thermal activation of dislocations became more operative than the hardening DSA effect for the given alloys tested at 500 ◦ C resulting in cyclic soft-
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Fig. 6. Monotonic stress–strain curves of 17-4 PH stainless steels in three conditions tested at 400 ◦ C.
ening. Moreover, the authors’ previous investigations [6,7] indicated that coarsening of Cu-rich phases was observed in Conditions A and H900 tested at 500 ◦ C and resulted in a reduction of hardness and fatigue resistance. Consequently, in addition to dynamic dislocation recovery, the coarsening of Cu-rich precipitates is believed to contribute to cyclic softening in Conditions A and H900 at 500 ◦ C. Furthermore, due to an increase in the period of a cycle with decreasing strain rate, the coarsening effect would become more effective such that the stress level was significantly decreased with
Fig. 8. TEM micrographs of microstructure in an H900 fatigue specimen tested at 500 ◦ C with a strain rate of 10−4 s−1 : (a) cell structures and (b) Cu-rich precipitates.
a reduction in strain rate for Conditions A and H900 tested at 500 ◦ C. 3.2. Effect of strain rate on fatigue life
Fig. 7. Planar slip bands formed in a Condition A fatigue specimen tested at 300 ◦ C with a strain rate of 10−4 s−1 .
Fig. 9 shows the variation of fatigue life as a function of strain rate at different temperatures for the heat-treated 17-4 PH stainless steels. It is seen in Fig. 9 that fatigue life of each condition, tested at a given temperature, decreased with a reduction in strain rate from 10−2 to 10−4 s−1 as a result of an enhanced DSA effect. DSA is considered to promote nonuniform fatigue deformation leading to localization of plastic flow and formation of intense slip bands [9,15]. In this regard, it would accelerate both crack initiation and
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Fig. 9. Variation of fatigue life with strain rate at various temperatures for 17-4 PH in different conditions: (a) Condition A, (b) H900, and (c) H1150.
propagation resulting in a significant reduction in fatigue life [9,15]. Moreover, under the DSA influence, the higher stress response developed during cyclic deformation caused a larger concentration of stress at the crack tip and an increase in crack growth rate leading to shorter crack propagation stage [9,11]. Generally, in the situations where DSA is present, the effect of DSA would be pronounced at decreasing strain rates or increasing temperatures [9]. Therefore, a decrease in strain rate from 10−2 to 10−4 s−1 , at a given temperature, makes the
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DSA effect far more effective to promote the development of planar slip bands and reduce the resistance to crack initiation and propagation leading to decreasing fatigue life. It has also been recognized by Srinivasan et al. [15] that a reduction in LCF life for a 316(N) stainless steel at 500–600 ◦ C, with decreasing strain rate, can be ascribed to the accelerated crack initiation and propagation caused by DSA. In addition, at 500 ◦ C and at a strain rate of 10−4 s−1 , a marked reduction in fatigue life of Conditions A and H900 in comparison with fatigue life at strain rates of 10−2 and 10−3 s−1 can be attributed to a coarsening of Cu-rich phases, which generates a degradation in material strength and greater plastic strain in each cycle. Fig. 9 indicates that for Conditions A and H900 the fatigue life in cycles, at a given strain rate, is not generally decreased with increasing temperature. In particular, at strain rates of 10−2 and 10−3 s−1 , fatigue life in either material condition has the highest value at the highest given temperature, 500 ◦ C. However, for the H1150 condition, the fatigue life at a given strain rate was generally independent of testing temperature between 300 and 500 ◦ C. It is believed that thermal dislocation recovery would induce greater material ductility and increase fatigue life under strain-control testing in terms of ductility exhaustion [16,17]. Therefore, in the current study, as testing temperature increased from 300 to 500 ◦ C, the enhanced thermal activation of dislocations is expected to increase material ductility and the stress needed to attain the imposed strain range decreases with increasing temperature. Hence, the fatigue lives at the higher temperatures can be effectively extended to a level close to or greater than at the lower temperature. An earlier study by Sandhya et al. [18] also indicated that LCF life of Alloy D9 increased with increasing temperature from 25 to 400 ◦ C due to reduction of stress needed to attain the applied strain range. However, the shorter fatigue lives of Conditions A and H900 tested at 500 ◦ C with a strain rate of 10−4 s−1 , as compared with the corresponding ones at 300 and 400 ◦ C, is attributed to the degradation of material strength and larger inelastic strain in each cycle caused by a pronounced coarsening effect of Curich precipitates. The total testing time was much longer at 10−4 s−1 strain rate such that coarsening effect of Cu-rich precipitates is more effective compared at the higher strain rates. SEM analyses indicated that at all given testing conditions, the fatigue fracture modes were commonly transgranular cracking accompanied by ductile striations on the fracture surface as shown in Fig. 10. Another example of transgranular cracking mode is given in Fig. 11 for a surface crack on the circumferential surface of a Condition H1150 specimen tested at 500 ◦ C and at a strain rate of 10−3 s−1 . In addition, no creep damages, such as cavities or intergranular microcracks, were observed for the given testing conditions. It has been reported that DSA would facilitate transgranular crack initiation and propagation in its operation regime [15]. Therefore, the crack initiation and propagation modes were transgranular at all of the given testing temperatures and stain rates.
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level followed by a rapid load drop due to formation of microcracks and their subsequent growth. 3. 17-4 PH alloys in the given three conditions at 500o C exhibited cyclic softening during LCF testing as a result of thermal dislocation recovery activities. Moreover, the cyclic softening behavior for Conditions A and H900 tested at 500o C is ascribed to the coarsening of Cu-rich precipitates. 4. For variously heat-treated 17-4 PH stainless steels tested at a given temperature, the number of cycles to failure consistently decreased with a reduction in strain rate due to an enhanced DSA effect. 5. At all given testing conditions, the fatigue failure for the given three material conditions occurred in a transgranular fracture mode.
Acknowledgement This work was funded by the National Science Council of the Republic of China (Taiwan) under Contract No. NSC-912216-E-008-007.
References Fig. 10. SEM fractographs of Condition H900 failed at 300 ◦ C with a strain rate of 10−3 s−1 : (a) initiation region and (b) stable crack growth region.
Fig. 11. SEM micrograph of a surface crack in a Condition H1150 fatigue specimen tested at 500 ◦ C with a strain rate of 10−3 s−1 .
4. Conclusions 1. The marked cyclic hardening behavior for Condition A tested at 300 and 400o C is attributed to dynamic strain aging effect during fatigue deformation. In addition, the in situ precipitation of Cu-rich phases is partially responsible for cyclic hardening in Condition A at 400o C. 2. The cyclic stress response of H900 and H1150 conditions tested at 300 and 400o C generally exhibited a stable stress
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