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Occurrence of dynamic strain aging in Alloy 617M under low cycle fatigue loading
Accepted Manuscript Occurrence of dynamic strain aging in Alloy 617M under low cycle fatigue loading Vani Shankar, Amit Kumar, K. Mariappan, R. Sandhya, K. Laha, A.K. Bhaduri, N. Narasaiah PII: DOI: Reference:
S0142-1123(17)30088-9 http://dx.doi.org/10.1016/j.ijfatigue.2017.03.001 JIJF 4264
To appear in:
International Journal of Fatigue
Received Date: Revised Date: Accepted Date:
22 November 2016 27 February 2017 1 March 2017
Please cite this article as: Shankar, V., Kumar, A., Mariappan, K., Sandhya, R., Laha, K., Bhaduri, A.K., Narasaiah, N., Occurrence of dynamic strain aging in Alloy 617M under low cycle fatigue loading, International Journal of Fatigue (2017), doi: http://dx.doi.org/10.1016/j.ijfatigue.2017.03.001
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Occurrence of dynamic strain aging in Alloy 617M under low cycle fatigue loading Vani Shankar*1, Amit Kumar2, K. Mariappan*, R. Sandhya*, K. Laha*, A. K. Bhaduri*and N. Narasaiah2 *Indira Gandhi Centre for Atomic Research, Kalpakkam-603102, India 2 Department of Metallurgical and Materials Engineering, National Institute of Technology Warangal-506 004, India India Abstract Low cycle fatigue (LCF) behavior of solution annealed Alloy 617M forging is studied at 300, 573, 773 and 973 K using strain amplitudes ±0.25, ±0.4, ±0.6 and ±0.8% at a nominal strain rate of 3×10 -3 s-1. The alloy evidenced the occurrence of dynamic strain aging (DSA) in the temperature regime 573 to 973 K caused by solute-dislocation interactions and the extent of interactions have been linked with the frequency and magnitude of stress drops/serrations on the hysteresis loops. The cyclic stress response, the hardening evolution curves, the frequency and magnitude of serrations on the hysteresis loops have been critically assessed in the framework of temperature-strain amplitude combination that affects the underlying fatigue deformation and the damage evolution process. At lower temperature-strain amplitude combination, the locking type B serrations appear from first cycle onwards that is persistent during the entire fatigue life. At 973 K, delayed strains for first serration to appear on hysteresis loop and disappearance of the type C serrations towards the later part of deformation have been associated with in-situ precipitation. Keywords: low cycle fatigue, Alloy 617M, dynamic strain aging, hysteresis loop Introduction Advanced Ultra Supercritical (A-USC) power plants are the most promising technology to improve thermal efficiency and reduce CO2 emissions of fossil fired power plants. A-USC power plants have working temperatures of approximately 993 K and a pressure of 35 MPa. In order to meet these working condition requirements, high strength superalloys are necessary. Alloy 617M is a candidate material for the components used in advanced ultra supercritical power plant. In A-USC power plants, the start-ups and shut-downs as well as power transients of the system creates low cycle fatigue (LCF) loading conditions on components such as critical steam piping, superheaters, reheaters, and steam turbine (blades, rotors, casing and propellers). Hence LCF is an important consideration in the design of systems operating at high temperature and subjected to thermal transients. At high temperatures the fatigue deformation and fatigue life is influenced by several temperature dependent mechanisms such as dynamic strain aging (DSA), oxidation, creep and phase transformations. These time dependent processes drastically reduce the fatigue life of the material [1]. Therefore, an awareness of the life limiting factors operative at service conditions is important for the safe operation of components. In order to develop a good understanding of the LCF behavior of Alloy 617M, experiments are carried out under various strain amplitude-temperature combinations. Through a detailed analysis of the hysteresis loops obtained over a range of temperature-strain amplitude combination, the underlying deformation and damage process operative under LCF condition has been identified. Literature review on Alloy 617 reveals that most of the work on the deformation behavior of this alloy is for the high temperature applications such as 1073 K and above since this alloy was a candidate material for components of high temperature gas cooled reactors. Not 1
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much work is reported on the deformation behavior of Alloy 617 in the temperature range of 300 to 973 K which also encompasses the DSA operative temperature regime (473 – 973 K) [2]. Hence the present work aims to understand the LCF behavior of Alloy 617M in the temperature range of 300 to 973 K. 2. EXPERIMENTAL The present study has been performed on forged product of Alloy 617M in the solution annealed condition. The chemical composition of the alloy is given in Table 1. Blanks of 22 x 22 x 110 mm were cut for machining of standard LCF specimens with uniform gauge section of 25 mm length and 10 mm diameter. Total axial strain controlled LCF experiments were performed in INSTRON servo-hydraulic system with electrical resistance furnace at strain amplitudes ranging from ± 0.25 to ± 0.8% and temperatures 300, 573, 773 and 973 K and strain rate, 3×10-3 s-1. The temperature was maintained within ±2 K over the entire gauge length. LCF testing was conducted in accordance with ASTM standard E606 [3]. Peak tensile stress at the half-life (i.e. at half of the number of cycles to failure) was taken as saturation or half-life stress and the cycle number corresponding to a drop of 20 % from the half-life stress was defined as the fatigue life. Optical and scanning electron microscopy was carried out to assess the changes in the microstructure during LCF experiments. Etching for optical microscopy was done on the polished surface of the untested and tested samples using Nital etchant with 90% methanol and 10% nitric acid. 3. RESULTS 3.1 Initial microstructure The initial microstructure of Alloy 617M forged product form is shown in Fig. 1. It consists of both large and small grains of austenite. The larger grains are about 60 μm and the smaller grains are about 10 to 15 μm. Linear intercept method was used to determine the grain size. The duplex distribution of grain size observed in the forged Alloy 617M has been ascribed to the fact that during solid solution annealing, the grains rich in precipitates exert a pinning stress on the grain boundaries and restrict the grain boundary movement (Zener pinning) there by preserving fine grained structure after annealing whereas the coarse grain is developed in the particle-free areas [4]. Annealing twins are also prevalent in the as-received specimens as confirmed by metallographic examinations of multiple specimens taken from various locations. 3.2 Effect of strain amplitude-temperature combination on shape of cyclic stress response (CSR) and hardening behavior Cyclic stress response (CSR) curves contain information about the underlying deformation and damage processes and hence in this section CSR curves have been critically analyzed to understand the deformation behavior as a function of strain amplitude and temperature combination. Cyclic stress response curves were constructed by plotting the peak tensile stress values (obtained from hysteresis loops) against the corresponding cycle number. The alloy is prone to cyclic hardening; there is an initial cyclic hardening period followed by either a brief saturation or softening until a sudden load drop occurs due to the development of macro crack leading to final failure (Figs. 2 and 3). Negative temperature dependence of cyclic stress has been observed in the temperature range 573 K to 973K and has been attributed to the increase in matrix hardening. The shape of the CSR (especially before the sudden load drop occurs due to macrocrack formation) is very sensitive to temperature-strain amplitude combination of the test (Figs. 2 and 3). The CSR has been categorized based on their shape as the ones showing (1) an initial brief hardening followed by cyclic softening, (2) a three slope behavior, i.e. initial brief sharp hardening followed by gradual hardening and then again sharp increase in the stress values, (3) a two slope behavior i.e. initial sharp increase followed by more 2
gradual increase in stress values and (4) a two slope behavior with initial gradual increase followed by sharper increase in the stress values in the latter part. Summary of the results are depicted in Table 2. Thus, at 300 K there is an initial brief hardening followed by cyclic softening, but at higher temperatures (573 K, 773 K and 973 K), the CSR show continuous cyclic hardening as the deformation progresses. Depending upon the temperature-strain amplitude combination, the CSR depicts one, two or even three slope hardening behavior. Thus at temperatures 573 K and 773 K (+ 0.25% and + 0.4%), a three slope behavior is observed whereas at the same temperatures for higher strain amplitudes (+ 0.6% and + 0.8%) and at 973 K for all the strain amplitudes, a two slope behavior is visible. A single slope behavior is observed for the highest strain amplitude + 0.8% and 773 K. Within the two slope behavior also the trends are varying. Hence whereas at lower temperatures (573 and 773 K), there is an initial sharp increase followed by a more gradual increase later, at higher temperatures (973 K), the trend is reversed i.e. the slope initially increases gradually which is taken over by a much sharper increase in the latter part. These observations point out that there definitely is a difference in the underlying deformation mechanism under different temperature-strain amplitude combination (i.e. at low (300 K), intermediate (573 K and 773 K) and high temperature (973 K)) which affects the overall cyclic stress response of Alloy 617M. An overall upward shift of the CSR curves with an increase in strain amplitude at all the four temperatures (300 K, 573 K, 773 K and 973 K) denotes that higher amount of stresses are required to sustain a higher amount of externally applied total strains that leads to a higher stress response. The amount of hardening (calculated from the cyclic stress response curve) incurred during cycling in various test conditions, is calculated with respect to the first cycle stress as given in equation (1). Amount of hardening =
(
N initial initial
)
(1)
where is the amount of hardening, initial is the peak tensile stress for first cycle and N is the peak tensile stress at a particular life fraction. The effect of varying strain amplitude (at constant temperature) or temperature (at constant strain amplitude) on the amount of hardening versus fatigue life (%) is depicted in Fig. 4(a) through (e). Same range of abscissa and ordinates are maintained for all the plots for making a reasonable comparison of the hardening behavior. It is noticed that the hardening behavior is sensitive to the applied strain amplitude and temperature combination. The hardening curves for 300 K for all the four strain amplitudes are nearly horizontal (unchanged with increasing fatigue life fractions) though there is an overall upward shift with an increase in the applied strain amplitude. As the temperature is increased to 573 K, the hardening evolution plot is no longer horizontal but increases linearly with increasing life fractions (Fig. 4 (b)). At temperatures 773 K and above, the hardening response is no longer linear but becomes parabolic. Also, the effect of increasing the strain amplitude on the overall upward shift of the hardening curves is substantial at 973 K. Figure 4 (e) makes a comparison of the hardening curves for the four temperatures at constant strain amplitude. It is observed that with an increase in temperature, the overall hardening curve shifts upwards (inverse temperature effect) and the shape of the hardening plot also changes. The hardening curve for 300 K is located at the lowest extremity and the 773 K and 973 K hardening curves are present at the upper extremity of the hardening versus fatigue life percentage plot. The change in the hardening evolution plot from linear at lower temperature (573 K) to a parabolic form at higher temperatures (773 K and 973 K) indicates a change in the underlying deformation mechanism and more than one hardening/softening factors being simultaneously 3
operative at these temperatures. The overall upward shift of the hardening evolution plot with increasing strain amplitude is also phenomenal at 773 K and above.
3.3 Hysteresis loop analysis Dynamic strain aging has been reported to occur in Alloy 617 in the temperature regime 473 – 973 K [2]. Dynamic strain aging (DSA) behavior under low cycle fatigue condition is not as well understood as under monotonic loading condition. The present study elucidates the trends followed by DSA under cyclic loading condition in terms of strength of the solute-dislocation interactions as a function of temperature-strain amplitude combination. Alloys which exhibit DSA often show one or more of the following deformation features during monotonic tensile tests; serrated flow, plateaus in the monotonic flow stress-temperature curves, high monotonic work hardening rates which increase with temperature, negative strain rate sensitivity of flow stress and ductility minima. It is seen that under LCF condition and some temperature-strain amplitude combination, the hysteresis loops depict serrations. At 300K, the hysteresis loops are smooth as no serrations are observed in the plastic portion of the hysteresis loops. Presence of serrations in the plastic strain region is one of the several manifestations for the occurrence of DSA and is caused by the temporary arrest of moving dislocations by the solute atmospheres present in the matrix [5-8]. Under LCF, apart from serrations in the plastic regions of hysteresis loops, DSA is also manifested in the form of increase in peak tensile stress with increasing temperature, increase in half-life stress with decreasing strain rate, increased cyclic hardening rate and decrease in plastic strain (obtained from the width of hysteresis loop at zero stress) [9, 10]. Under monotonic loading conditions, solute species causing DSA are identified by utilizing the wellestablished criterion of calculating minimum plastic strain for the appearance of first serration and use of empirical relationships to calculate activation energy for solute diffusion [1, 11,12]. This indicates that a minimum number of mobile dislocations should be present and the velocity of moving dislocations (strain rate) and diffusing solutes (temperature) should be optimum and matching for the dislocation-solute interactions to occur. In similar lines, serrations have been observed after certain amount of plastic strains in the hysteresis loops. It is also observed that the strains for the first serration to appear in subsequent cycles are not the same. Additionally, depending upon the temperature-strain amplitude combination, the serration characteristics (in terms of frequency and magnitude) also varies. As mentioned earlier, there are no serrations in the hysteresis loops for the tests carried out at 300K (Fig. 5 (a)), whereas at temperature-strain amplitude combination (at 573 K and above), serrations appear in the hysteresis loops. Typical representations are made in Fig. 5 (b) and (c). Serrations appear on the plastic portion of the hysteresis loops due to the occurrence of DSA; stresses rise because the moving dislocations are temporarily locked in by the solute atmospheres and higher stresses are required to break free from the interaction atmosphere until relocking of the freed dislocations by another solute atmosphere occurs. Higher is the dislocation-solute interaction strength, greater shall be the stresses required to break free the interaction field and hence larger shall be the stress drop/serration height. In the present study it is seen that the serration characteristics (i.e the frequency of appearance of the serrations on hysteresis loops and the magnitude of stress drops on the hysteresis loops) are greatly governed by the temperature-strain amplitude combination. 3.4 Effect of strain amplitude-temperature combination on evolution of strength/magnitude of DSA interaction The amount of stress drops and the strain for the first serration to appear in a hysteresis loop as a function of test condition have been determined. It is found that at 573 K, equally spaced serrations with magnitude 10-15 MPa appear right from first cycle onwards and that is persistent until failure. At 773 K, the serrations appear 4
at regular intervals also but are of much larger magnitudes; initially there are nearly 60 MPa stress drops that diminish down to ~25 MPa after many cycles. At 973 K, only the initial 7 cycles show large stress drops of the order of 125 MPa and then the serrations disappear. It is also interesting to note that at this temperature, whereas large stress drops appear in the tensile portion of hysteresis loops for the first initial cycles, the compressive portion of the hysteresis loop contains equally spaced serrations of comparatively very small magnitudes thereby indicating that the strength of interactions in tensile and compressive portions of the hysteresis loops are not the same and is dependent on the direction of straining. Under all the test conditions where serrations appear on hysteresis loops, the magnitude of serrations does not remain the same throughout the fatigue life; in the initial cycles larger magnitude serrations are taken over by the comparatively lower magnitude serrations in the latter part. Strength/ magnitude of DSA/solute-dislocation interactions are depicted in Figure 6 in terms of the magnitude of the largest serration in a given cycle versus the cycle number (Fig. 6 (a)) and the strain (total) at which the first serration appear in a given cycle (Fig. 6 (b)). As stated above and depicted in Fig. 6 (a), the magnitude of serration decreases and attains a saturation value with the progress of deformation (increasing number of cycles). Strain at which the first serration appear in a hysteresis loop also increases with the progress of deformation (Fig. 6(b)). This holds true for all the four strain amplitudes. The effect of temperature on the serration magnitude and strain for the first serration to appear is shown in Figs. 6 (c) and (d) respectively. The plot clearly indicates a strong temperature dependence of the magnitude of stress drops. It is also interesting to note that at 773K, the minimum strain required for the first serration to appear (Fig. 6b), is the lowest among the three temperatures (i.e. 573, 773 and 973 K) thereby indicating that the occurrence of DSA is the most favored at 773 K than at any other temperatures. The strength of the solute-dislocation interaction (in terms of the magnitude of serrations) at 773 K is also significantly larger than those observed at 573 K. Another noteworthy feature in all the four plots are that both maximum stress drop in a cycle and the strain to first serration show large changes during the initial cycles only and which is then followed by a saturation period (Fig. 6). The appearance of serrations on hysteresis loops have been defined in the framework of the well-defined classification of serrations [1, 13, 14.]. Thus at lower temperatures such as at 573 K, type B serrations appear and the DSA is caused by the solutes interactions with the dislocations. Type B serrations are known to fluctuate about the hysteresis loop and are also considered as ‘‘locking type serrations” [15]. At temperatures 773 K and above, type C serrations of unlocking nature are responsible for the serrated behavior. The type C serrations have been reported to occur in alloys at high temperatures [1, 7]. At 973 K though the stress drop is very large (125 MPa) as compared to other two temperatures (573 K and 773 K) (Fig. 6 (c)), the initial strain required for the first serration to appear is also higher as compared to 573 and 773 K (Fig. 6 (d)). Such delayed serrations and disappearance of serrations towards the later part of deformation have been associated with in-situ precipitation in other nickel base superalloys such as Alloy 625 [16].The highest magnitude of stress drops (~ 125 MPa at 973 K) for initial few cycles indicate strongest interaction of solute-dislocation atmospheres and then the solutes are no longer freely available to interact with the moving dislocations. This happens due to precipitation and is manifested in the form of disappearance of the serrations. 4.0 DISCUSSION 4.1 Cyclic stress response The inverse temperature dependence of CSR has been observed in the temperature regime 573 to 973 K (identified as the DSA region) due to matrix hardening. At room temperature (300 K) the hardening is caused by dislocation multiplication alone due to cyclic deformation. On increasing the temperature from 300 K to 5
973 K an additional amount of matrix hardening occur due to DSA. It is an attractive interaction between the solute species and mobile dislocations [8, 17] during deformation, and thus it constrains the dislocation motion to the slip plane. Consequently, in order to maintain an imposed strain rate, cyclic flow stress increases either to unlock the dislocations from the obstacles or to generate new dislocations. The moving dislocations get momentarily locked around the elastic strain fields of solute particles present in the matrix [1, 8] causing additional hardening of the matrix. At temperatures, where the effects of time dependent processes such as creep and oxidation are found to be minimal, the drastic reduction in LCF life has been reported to occur due to the deleterious effects of DSA in other alloys as well such as Nimonic PE 16 superalloy [9], Haynes 188 superalloy [10], Hastealloy [18] and Duplex stainless steel [19], 316 L(N) [20]. The cyclic hardening in the initial cycles observed in this alloy has been ascribed to the dislocation multiplication and accumulation within the slip band [21]. At elevated temperatures in addition to the dislocation multiplication and accumulation within the slip band [21] the continuous hardening is due to the formation of dislocation substructure which provides sites for numerous heterogeneous precipitations of fine M 23C6 particles. This in turn stabilizes the substructure during further cycling and prevents the recovery processes [22]. Srinivasan et al. [23] have observed an increase in friction stress with increasing temperature in the DSA regime and attributed to the interaction of solute atoms with mobile dislocations. The mechanism of the one/two/three slope behavior of CSR can be explained based on balance of factors affecting the deformation behavior. At ambient temperature, only dislocation multiplication occurs, whereas at elevated temperatures, the availability or non-availability of solutes for the interactions of solutes and dislocations decides the overall cyclic stress response. Detailed microstructural studies by Kaoumi et al. [24] recently confirmed that precipitate evolution occurs above 873 K, i.e. dissolution of the original primary carbides (MC or M6C formed during solidification of ingots) adds a significant amount of solute atoms into the matrix which act as barriers to dislocation motion, and leads to re-precipitation of smaller M23C6 carbides. At 1223 K, strengthening provided by dissolution of original primary carbides (MC or M6C) and reprecipitation of smaller M23C6 carbides competes with the softening process of recrystallization. The primary MC carbides present in the undissolved state during solution annealing have been reported to decompose into M23C6 and M6C on prolonged exposure at elevated temperatures [25]. At higher temperatures such as 973 K and above, a possibility of in-situ precipitation cannot be ruled out as has been reported to occur in Alloy 625 [16]. In Alloy 617, Wu et al. [26] indicated 973 K as the nose of the time-temperature-transformation diagram that favors formation of Ti(C, N), M23C23 and γ ׳phases. Maier et al. [27] recently reported that mere thermal aging of Alloy 617 at 973 K did not result in any significant microstructural changes, but under fatigue loading, significant amount of M 23C6 carbides precipitated within a relatively short time. The disappearance of serrations from hysteresis loops after few cycles at 973 K is associated with the unavailability of solute particles due to in-situ precipitate formation during the initial periods. Disappearance of serrations from flow curves have been also reported by Hayes [28] and Hayes and Hayes [29] in alloys like AISI 1020, 2.25 Cr-1Mo, alloy 718, and alloy 600 and by Vani Shankar et al. in Alloy 625 [16] during monotonic tensile tests. It has been suggested [30] that unlocking serrations are connected with precipitation before and during the test. Hayes [28] and Hayes and Hayes [29] carried out systematic investigations in AISI 1020, 2.25 Cr-1Mo, alloy 718, and alloy 600 to relate the strain for disappearance of serrations with strain rate and temperature. They have shown that the disappearance of serrations from the flow curve occurs in the high-temperature regime by either a progressively longer strain to the onset of serrations (a critical-strain-delay mechanism) or by a progressively smaller strain to the disappearance of serrations (disappearance off the end of the flow curve). In the former case, carbon diffusing down the dislocation line to a precipitate sink is suggested as the mechanism; however, in the latter case, carbon reacting with a carbide forming species on the dislocation line is responsible for the disappearance of 6
serrations. Thus, the important conclusion from their studies is that the disappearance is generally related to a precipitation mechanism, and the disappearance of serrated flow would occur when a balance is reached between the growth of carbon atmosphere and its depletion due to reaction between substitutional (carbideforming) atoms in the atmosphere. In this context, it has been suggested that the occurrence of type-C serrations could be regarded as a precursor to the precipitation and disappearance of serrations [30]. 4.2 Evolution of hysteresis loops and serrations: strength/magnitude of DSA interaction As stated earlier, the hysteresis loops showed serrations/stress drops in the plastic strained regions whose frequency of appearance and magnitude depended upon temperature-strain amplitude combination. Also, the serration heights varied during the life span of fatigue cycling. Since additional stresses are required to break free the solute-dislocation interaction atmospheres, the strength/magnitude of the interactions is evaluated by the magnitude of stress drops. The minimum strain required for first serration to appear in a hysteresis loop symbolizes how favorable it is to cause solute-dislocation interaction (DSA). Stress drop magnitude can be also expected to depend upon the solute species responsible for causing DSA. It is seen that the stress drops are significantly large during the initial few cycles and then the values saturates (Fig. 6 (a)). This denotes stronger interactions of solutes and dislocations during the initial periods of cyclic deformation and then formation of a stable microstructure/ solute-dislocation interaction zone. Once dislocations are generated during the initial deformation in a cycle, the diffusing solutes can interact with the moving dislocations and can cause strong solute-dislocation interactions. As the cyclic deformation progresses, the dislocations form a stable/constant microstructure in which the solutes come across almost constant number of dislocation atmospheres which results in stable stress or strain values after some time. Microstructural based quantification is however necessary to confirm the aforesaid. At 773 K, the magnitude of serrations is comparatively higher (among 573 and 773 K), and the minimum strain for first serration to appear is the lowest (among three temperatures, 573, 773 and 973 K). This accounts for the most favorable temperature for the occurrence of DSA at the strain rate of the LCF tests. The disappearance of serrations off the hysteresis loops after initial few cycles (Fig. 5 (c)) and delayed strains for the appearance of first serration at 973 K indicates that a different mechanism such as in-situ precipitation is operative. Numerous vacancies are generated during cyclic deformation that enhances in-situ precipitation. Another noteworthy feature is that there is a significant difference between the magnitude of serrations observed at 973 K than those at lower temperatures. Thus at 973 K the magnitude of serration is exceedingly higher (~ 125 MPa) than those observed at lower temperatures such as 773 K (10 to 60 MPa maximum). This amounts to the fact that the species responsible for causing DSA might be different at elevated temperatures such as 973 K than those operative at lower temperatures (below 973 K). Two temperature-strain rate combination regions in which two different species are responsible to cause DSA have been identified in a similar alloy system such as Alloy 625 under monotonic loading condition as well [16]. Thus whereas interstitial solutes such as carbon were responsible to cause DSA at lower temperatures, substitutional solute such as molybdenum caused DSA at higher temperatures in Alloy 625. It was stated that solutes were unavailable to cause further dislocation-solute interactions due to precipitation in Alloy 625 and hence serrations disappeared from the tensile flow curves after some time. Recently Ekaputra et al. [2] under monotonic loading condition has identified three zones for the occurrence of DSA in Alloy 617. Through activation energy calculation for solute diffusion in the nickel matrix, three temperature zones were identified for causing DSA. Thus Type D serrations occurred at 473 K (activation energy of 77.8 kJ/mol corresponded to pipe diffusion of carbon), Type A+B serrations between 573 - 773 K (activation energy ranged from 110.7 kJ/mol to 193.8 kJ/mol and was attributed to lattice diffusion of interstitial solute atoms of carbon) and Type C serrations that appeared between 873 to 973 K (activation energy of 265.1 kJ/mol, corresponded to the 7
lattice diffusion of substitutional solute atoms, such as chromium). In similar lines it may be stated that under cyclic loading conditions also, while carbon solutes causes DSA at lower temperatures, that at higher temperatures, Mo or Cr might be responsible. In-situ precipitation of M23C6 causes disappearance of serrations off the hysteresis loops after few cycles at 973 K. Detailed microscopic examination and determination of activation energy for the solutes responsible for causing DSA under cyclic loading is however necessary to confirm the aforesaid.
CONCLUSIONS 1) Occurrence of dynamic strain aging (DSA) in the temperature range 573 to 973 K is the cause for matrix hardening in Alloy 617M. 2) The magnitude and frequency of appearance of the serrations in hysteresis loops are linked with the strength of the solute-dislocation interactions causing DSA and is sensitive to strain amplitudetemperature combination. 3) 773 K is the most favored temperature for the occurrence of DSA in Alloy 617M in the range of temperatures and strain rate employed in this study. Two temperature regimes in which two species might be responsible to cause DSA in this alloy exists. At temperatures below 973 K, B type serrations are caused by interstitial carbon solutes whereas at 973 K other substitutional elements are responsible to cause C-type serrations and DSA. Appearance and then disappearance of the serrations after few cycles at 973 K indicate unavailability of the carbide forming element in the matrix possibly due to in-situ precipitation of M23C6.
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[17] A.W. Sleeswyk, Acta Metall6 (9)(1958)598–603. [18] M. G. Castelli, R. V. Miner and D. N. Robinson, ASTM-STP 1186, ASTM, Philadelphia, 1993;106 [19] S. Herenu, I. Alvarez-Armas and A. F. Armas, Scr. Mater. 45 (2001) 739-745. [20] V. S. Srinivasan, M. Valsan, R. Sandhya, K. Bhanu Sankara Rao, S. L. Mannan and D. H. Sastry, Int. J. Fatigue 21 (1999) 11-21. [21] G. Chai, P. Liu and J. Frodigh, J. Mater Sci., 39 (2004) 2689-2697. [22] M. A. Burke, and C. G. Beck, Met Trans A 15A (1984) 661-670. [23] V.S. Srinivasan, R. Sandhya, M. Valsan, K. Bhanu Sankara Rao, S.L. Mannan and D.H. Sastry , Scripta Materialia, 37(10)(1997) 1593-l 598. [24] D. Kaoumi, K. Hrutkay, Journal of Nuclear Materials 454 (2014) 265–273. [25] D.R. Muzyka: in The Superalloys 718, C.T. Sims and W.C. Hagel, eds., Wiley, New York, NY, 1972, p. 113. [26] Q. Wu, H. Song, R. W. Swindeman, Metallurgical and Materials Transactions 39A(11) (2008) 25692585. [27] Gerhard Maier, Hermann Riedel, Christoph Somsen, International Journal of Fatigue 55 (2013) 126–135. [28]R.W. Hayes, Acta Metall., 31 (1983) 365-371. [29] R.W. Hayes and W.C. Hayes: Acta Metall., 32(1984)259-267. [30] S. Venkadesan, C. Phaniraj, P.V. Sivaprasad, and P. Rodriguez, Acta Metall., 40 (1992)569-580.
Figure 1 Optical micrograph depicting initial microstructure of untested sample
9
(a)
(b) Figure 2 Effect of strain amplitude on the cyclic stress response of Alloy 617M at (a) 300 K and 773 K.
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Figure 3 Effect of temperature on the cyclic stress response of Alloy 617M at +0.6% strain amplitude.
(a)
(b)
(c)
(d)
11
(e) Figure 4 Effect of change in strain amplitude (at constant temperature such as at (a) 300 K, (b) 573 K, (c) 773 K and (d) 973 K) and (e) temperature (at constant strain amplitude, +0.4%) on the cyclic hardening evolution curve of Alloy 617 M.
(a)
(b)
(c) Figure 5 Representation of the evolution of hysteresis loops showing effects of temperature- through tests carried out at (a) 300 K; +0.6%, (b) 773 K; +0.6% and (c) 973 K; +0.6%. 12
(a)
(b)
(c) (d) Fig. 6 Plot of maximum stress drop in a cycle ((a) and (c)) and critical strain for appearance of first serration in hysteresis loop ((b) and (d)) versus the number of cycles showing for various strain amplitudes ((a) and (b)) and temperatures ((c) and (d)).
13
Table 1 Chemical composition of Alloy 617M forging Element
C
Mo
Fe
Co
Ti
Cr
Si
Al
B
Ni
Wt% (F)
0.07
9.38
0.14
12.02
0.42
22.12
0.02
1.12
0.004
Bal.
Table 2 Summary of shape of cyclic stress response curve as a function of temperature-strain amplitude combination of the test
14
Highlights 1) Dynamic strain aging (DSA) occurs between 573 and 973 K and causes matrix hardening. 2) Magnitude and frequency of appearance of serrations in hysteresis loops linked with strength of solutedislocation interactions. 3) 773 K is the most favored temperature for occurrence of DSA. 4) More than one species responsible for causing DSA.
15
S0142-1123(17)30088-9 http://dx.doi.org/10.1016/j.ijfatigue.2017.03.001 JIJF 4264
To appear in:
International Journal of Fatigue
Received Date: Revised Date: Accepted Date:
22 November 2016 27 February 2017 1 March 2017
Please cite this article as: Shankar, V., Kumar, A., Mariappan, K., Sandhya, R., Laha, K., Bhaduri, A.K., Narasaiah, N., Occurrence of dynamic strain aging in Alloy 617M under low cycle fatigue loading, International Journal of Fatigue (2017), doi: http://dx.doi.org/10.1016/j.ijfatigue.2017.03.001
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Occurrence of dynamic strain aging in Alloy 617M under low cycle fatigue loading Vani Shankar*1, Amit Kumar2, K. Mariappan*, R. Sandhya*, K. Laha*, A. K. Bhaduri*and N. Narasaiah2 *Indira Gandhi Centre for Atomic Research, Kalpakkam-603102, India 2 Department of Metallurgical and Materials Engineering, National Institute of Technology Warangal-506 004, India India Abstract Low cycle fatigue (LCF) behavior of solution annealed Alloy 617M forging is studied at 300, 573, 773 and 973 K using strain amplitudes ±0.25, ±0.4, ±0.6 and ±0.8% at a nominal strain rate of 3×10 -3 s-1. The alloy evidenced the occurrence of dynamic strain aging (DSA) in the temperature regime 573 to 973 K caused by solute-dislocation interactions and the extent of interactions have been linked with the frequency and magnitude of stress drops/serrations on the hysteresis loops. The cyclic stress response, the hardening evolution curves, the frequency and magnitude of serrations on the hysteresis loops have been critically assessed in the framework of temperature-strain amplitude combination that affects the underlying fatigue deformation and the damage evolution process. At lower temperature-strain amplitude combination, the locking type B serrations appear from first cycle onwards that is persistent during the entire fatigue life. At 973 K, delayed strains for first serration to appear on hysteresis loop and disappearance of the type C serrations towards the later part of deformation have been associated with in-situ precipitation. Keywords: low cycle fatigue, Alloy 617M, dynamic strain aging, hysteresis loop Introduction Advanced Ultra Supercritical (A-USC) power plants are the most promising technology to improve thermal efficiency and reduce CO2 emissions of fossil fired power plants. A-USC power plants have working temperatures of approximately 993 K and a pressure of 35 MPa. In order to meet these working condition requirements, high strength superalloys are necessary. Alloy 617M is a candidate material for the components used in advanced ultra supercritical power plant. In A-USC power plants, the start-ups and shut-downs as well as power transients of the system creates low cycle fatigue (LCF) loading conditions on components such as critical steam piping, superheaters, reheaters, and steam turbine (blades, rotors, casing and propellers). Hence LCF is an important consideration in the design of systems operating at high temperature and subjected to thermal transients. At high temperatures the fatigue deformation and fatigue life is influenced by several temperature dependent mechanisms such as dynamic strain aging (DSA), oxidation, creep and phase transformations. These time dependent processes drastically reduce the fatigue life of the material [1]. Therefore, an awareness of the life limiting factors operative at service conditions is important for the safe operation of components. In order to develop a good understanding of the LCF behavior of Alloy 617M, experiments are carried out under various strain amplitude-temperature combinations. Through a detailed analysis of the hysteresis loops obtained over a range of temperature-strain amplitude combination, the underlying deformation and damage process operative under LCF condition has been identified. Literature review on Alloy 617 reveals that most of the work on the deformation behavior of this alloy is for the high temperature applications such as 1073 K and above since this alloy was a candidate material for components of high temperature gas cooled reactors. Not 1
Corresponding author: Tel. +91-44-274800-ext21147 Fax, +91-44-27480075 E-mail: [email protected] 1
much work is reported on the deformation behavior of Alloy 617 in the temperature range of 300 to 973 K which also encompasses the DSA operative temperature regime (473 – 973 K) [2]. Hence the present work aims to understand the LCF behavior of Alloy 617M in the temperature range of 300 to 973 K. 2. EXPERIMENTAL The present study has been performed on forged product of Alloy 617M in the solution annealed condition. The chemical composition of the alloy is given in Table 1. Blanks of 22 x 22 x 110 mm were cut for machining of standard LCF specimens with uniform gauge section of 25 mm length and 10 mm diameter. Total axial strain controlled LCF experiments were performed in INSTRON servo-hydraulic system with electrical resistance furnace at strain amplitudes ranging from ± 0.25 to ± 0.8% and temperatures 300, 573, 773 and 973 K and strain rate, 3×10-3 s-1. The temperature was maintained within ±2 K over the entire gauge length. LCF testing was conducted in accordance with ASTM standard E606 [3]. Peak tensile stress at the half-life (i.e. at half of the number of cycles to failure) was taken as saturation or half-life stress and the cycle number corresponding to a drop of 20 % from the half-life stress was defined as the fatigue life. Optical and scanning electron microscopy was carried out to assess the changes in the microstructure during LCF experiments. Etching for optical microscopy was done on the polished surface of the untested and tested samples using Nital etchant with 90% methanol and 10% nitric acid. 3. RESULTS 3.1 Initial microstructure The initial microstructure of Alloy 617M forged product form is shown in Fig. 1. It consists of both large and small grains of austenite. The larger grains are about 60 μm and the smaller grains are about 10 to 15 μm. Linear intercept method was used to determine the grain size. The duplex distribution of grain size observed in the forged Alloy 617M has been ascribed to the fact that during solid solution annealing, the grains rich in precipitates exert a pinning stress on the grain boundaries and restrict the grain boundary movement (Zener pinning) there by preserving fine grained structure after annealing whereas the coarse grain is developed in the particle-free areas [4]. Annealing twins are also prevalent in the as-received specimens as confirmed by metallographic examinations of multiple specimens taken from various locations. 3.2 Effect of strain amplitude-temperature combination on shape of cyclic stress response (CSR) and hardening behavior Cyclic stress response (CSR) curves contain information about the underlying deformation and damage processes and hence in this section CSR curves have been critically analyzed to understand the deformation behavior as a function of strain amplitude and temperature combination. Cyclic stress response curves were constructed by plotting the peak tensile stress values (obtained from hysteresis loops) against the corresponding cycle number. The alloy is prone to cyclic hardening; there is an initial cyclic hardening period followed by either a brief saturation or softening until a sudden load drop occurs due to the development of macro crack leading to final failure (Figs. 2 and 3). Negative temperature dependence of cyclic stress has been observed in the temperature range 573 K to 973K and has been attributed to the increase in matrix hardening. The shape of the CSR (especially before the sudden load drop occurs due to macrocrack formation) is very sensitive to temperature-strain amplitude combination of the test (Figs. 2 and 3). The CSR has been categorized based on their shape as the ones showing (1) an initial brief hardening followed by cyclic softening, (2) a three slope behavior, i.e. initial brief sharp hardening followed by gradual hardening and then again sharp increase in the stress values, (3) a two slope behavior i.e. initial sharp increase followed by more 2
gradual increase in stress values and (4) a two slope behavior with initial gradual increase followed by sharper increase in the stress values in the latter part. Summary of the results are depicted in Table 2. Thus, at 300 K there is an initial brief hardening followed by cyclic softening, but at higher temperatures (573 K, 773 K and 973 K), the CSR show continuous cyclic hardening as the deformation progresses. Depending upon the temperature-strain amplitude combination, the CSR depicts one, two or even three slope hardening behavior. Thus at temperatures 573 K and 773 K (+ 0.25% and + 0.4%), a three slope behavior is observed whereas at the same temperatures for higher strain amplitudes (+ 0.6% and + 0.8%) and at 973 K for all the strain amplitudes, a two slope behavior is visible. A single slope behavior is observed for the highest strain amplitude + 0.8% and 773 K. Within the two slope behavior also the trends are varying. Hence whereas at lower temperatures (573 and 773 K), there is an initial sharp increase followed by a more gradual increase later, at higher temperatures (973 K), the trend is reversed i.e. the slope initially increases gradually which is taken over by a much sharper increase in the latter part. These observations point out that there definitely is a difference in the underlying deformation mechanism under different temperature-strain amplitude combination (i.e. at low (300 K), intermediate (573 K and 773 K) and high temperature (973 K)) which affects the overall cyclic stress response of Alloy 617M. An overall upward shift of the CSR curves with an increase in strain amplitude at all the four temperatures (300 K, 573 K, 773 K and 973 K) denotes that higher amount of stresses are required to sustain a higher amount of externally applied total strains that leads to a higher stress response. The amount of hardening (calculated from the cyclic stress response curve) incurred during cycling in various test conditions, is calculated with respect to the first cycle stress as given in equation (1). Amount of hardening =
(
N initial initial
)
(1)
where is the amount of hardening, initial is the peak tensile stress for first cycle and N is the peak tensile stress at a particular life fraction. The effect of varying strain amplitude (at constant temperature) or temperature (at constant strain amplitude) on the amount of hardening versus fatigue life (%) is depicted in Fig. 4(a) through (e). Same range of abscissa and ordinates are maintained for all the plots for making a reasonable comparison of the hardening behavior. It is noticed that the hardening behavior is sensitive to the applied strain amplitude and temperature combination. The hardening curves for 300 K for all the four strain amplitudes are nearly horizontal (unchanged with increasing fatigue life fractions) though there is an overall upward shift with an increase in the applied strain amplitude. As the temperature is increased to 573 K, the hardening evolution plot is no longer horizontal but increases linearly with increasing life fractions (Fig. 4 (b)). At temperatures 773 K and above, the hardening response is no longer linear but becomes parabolic. Also, the effect of increasing the strain amplitude on the overall upward shift of the hardening curves is substantial at 973 K. Figure 4 (e) makes a comparison of the hardening curves for the four temperatures at constant strain amplitude. It is observed that with an increase in temperature, the overall hardening curve shifts upwards (inverse temperature effect) and the shape of the hardening plot also changes. The hardening curve for 300 K is located at the lowest extremity and the 773 K and 973 K hardening curves are present at the upper extremity of the hardening versus fatigue life percentage plot. The change in the hardening evolution plot from linear at lower temperature (573 K) to a parabolic form at higher temperatures (773 K and 973 K) indicates a change in the underlying deformation mechanism and more than one hardening/softening factors being simultaneously 3
operative at these temperatures. The overall upward shift of the hardening evolution plot with increasing strain amplitude is also phenomenal at 773 K and above.
3.3 Hysteresis loop analysis Dynamic strain aging has been reported to occur in Alloy 617 in the temperature regime 473 – 973 K [2]. Dynamic strain aging (DSA) behavior under low cycle fatigue condition is not as well understood as under monotonic loading condition. The present study elucidates the trends followed by DSA under cyclic loading condition in terms of strength of the solute-dislocation interactions as a function of temperature-strain amplitude combination. Alloys which exhibit DSA often show one or more of the following deformation features during monotonic tensile tests; serrated flow, plateaus in the monotonic flow stress-temperature curves, high monotonic work hardening rates which increase with temperature, negative strain rate sensitivity of flow stress and ductility minima. It is seen that under LCF condition and some temperature-strain amplitude combination, the hysteresis loops depict serrations. At 300K, the hysteresis loops are smooth as no serrations are observed in the plastic portion of the hysteresis loops. Presence of serrations in the plastic strain region is one of the several manifestations for the occurrence of DSA and is caused by the temporary arrest of moving dislocations by the solute atmospheres present in the matrix [5-8]. Under LCF, apart from serrations in the plastic regions of hysteresis loops, DSA is also manifested in the form of increase in peak tensile stress with increasing temperature, increase in half-life stress with decreasing strain rate, increased cyclic hardening rate and decrease in plastic strain (obtained from the width of hysteresis loop at zero stress) [9, 10]. Under monotonic loading conditions, solute species causing DSA are identified by utilizing the wellestablished criterion of calculating minimum plastic strain for the appearance of first serration and use of empirical relationships to calculate activation energy for solute diffusion [1, 11,12]. This indicates that a minimum number of mobile dislocations should be present and the velocity of moving dislocations (strain rate) and diffusing solutes (temperature) should be optimum and matching for the dislocation-solute interactions to occur. In similar lines, serrations have been observed after certain amount of plastic strains in the hysteresis loops. It is also observed that the strains for the first serration to appear in subsequent cycles are not the same. Additionally, depending upon the temperature-strain amplitude combination, the serration characteristics (in terms of frequency and magnitude) also varies. As mentioned earlier, there are no serrations in the hysteresis loops for the tests carried out at 300K (Fig. 5 (a)), whereas at temperature-strain amplitude combination (at 573 K and above), serrations appear in the hysteresis loops. Typical representations are made in Fig. 5 (b) and (c). Serrations appear on the plastic portion of the hysteresis loops due to the occurrence of DSA; stresses rise because the moving dislocations are temporarily locked in by the solute atmospheres and higher stresses are required to break free from the interaction atmosphere until relocking of the freed dislocations by another solute atmosphere occurs. Higher is the dislocation-solute interaction strength, greater shall be the stresses required to break free the interaction field and hence larger shall be the stress drop/serration height. In the present study it is seen that the serration characteristics (i.e the frequency of appearance of the serrations on hysteresis loops and the magnitude of stress drops on the hysteresis loops) are greatly governed by the temperature-strain amplitude combination. 3.4 Effect of strain amplitude-temperature combination on evolution of strength/magnitude of DSA interaction The amount of stress drops and the strain for the first serration to appear in a hysteresis loop as a function of test condition have been determined. It is found that at 573 K, equally spaced serrations with magnitude 10-15 MPa appear right from first cycle onwards and that is persistent until failure. At 773 K, the serrations appear 4
at regular intervals also but are of much larger magnitudes; initially there are nearly 60 MPa stress drops that diminish down to ~25 MPa after many cycles. At 973 K, only the initial 7 cycles show large stress drops of the order of 125 MPa and then the serrations disappear. It is also interesting to note that at this temperature, whereas large stress drops appear in the tensile portion of hysteresis loops for the first initial cycles, the compressive portion of the hysteresis loop contains equally spaced serrations of comparatively very small magnitudes thereby indicating that the strength of interactions in tensile and compressive portions of the hysteresis loops are not the same and is dependent on the direction of straining. Under all the test conditions where serrations appear on hysteresis loops, the magnitude of serrations does not remain the same throughout the fatigue life; in the initial cycles larger magnitude serrations are taken over by the comparatively lower magnitude serrations in the latter part. Strength/ magnitude of DSA/solute-dislocation interactions are depicted in Figure 6 in terms of the magnitude of the largest serration in a given cycle versus the cycle number (Fig. 6 (a)) and the strain (total) at which the first serration appear in a given cycle (Fig. 6 (b)). As stated above and depicted in Fig. 6 (a), the magnitude of serration decreases and attains a saturation value with the progress of deformation (increasing number of cycles). Strain at which the first serration appear in a hysteresis loop also increases with the progress of deformation (Fig. 6(b)). This holds true for all the four strain amplitudes. The effect of temperature on the serration magnitude and strain for the first serration to appear is shown in Figs. 6 (c) and (d) respectively. The plot clearly indicates a strong temperature dependence of the magnitude of stress drops. It is also interesting to note that at 773K, the minimum strain required for the first serration to appear (Fig. 6b), is the lowest among the three temperatures (i.e. 573, 773 and 973 K) thereby indicating that the occurrence of DSA is the most favored at 773 K than at any other temperatures. The strength of the solute-dislocation interaction (in terms of the magnitude of serrations) at 773 K is also significantly larger than those observed at 573 K. Another noteworthy feature in all the four plots are that both maximum stress drop in a cycle and the strain to first serration show large changes during the initial cycles only and which is then followed by a saturation period (Fig. 6). The appearance of serrations on hysteresis loops have been defined in the framework of the well-defined classification of serrations [1, 13, 14.]. Thus at lower temperatures such as at 573 K, type B serrations appear and the DSA is caused by the solutes interactions with the dislocations. Type B serrations are known to fluctuate about the hysteresis loop and are also considered as ‘‘locking type serrations” [15]. At temperatures 773 K and above, type C serrations of unlocking nature are responsible for the serrated behavior. The type C serrations have been reported to occur in alloys at high temperatures [1, 7]. At 973 K though the stress drop is very large (125 MPa) as compared to other two temperatures (573 K and 773 K) (Fig. 6 (c)), the initial strain required for the first serration to appear is also higher as compared to 573 and 773 K (Fig. 6 (d)). Such delayed serrations and disappearance of serrations towards the later part of deformation have been associated with in-situ precipitation in other nickel base superalloys such as Alloy 625 [16].The highest magnitude of stress drops (~ 125 MPa at 973 K) for initial few cycles indicate strongest interaction of solute-dislocation atmospheres and then the solutes are no longer freely available to interact with the moving dislocations. This happens due to precipitation and is manifested in the form of disappearance of the serrations. 4.0 DISCUSSION 4.1 Cyclic stress response The inverse temperature dependence of CSR has been observed in the temperature regime 573 to 973 K (identified as the DSA region) due to matrix hardening. At room temperature (300 K) the hardening is caused by dislocation multiplication alone due to cyclic deformation. On increasing the temperature from 300 K to 5
973 K an additional amount of matrix hardening occur due to DSA. It is an attractive interaction between the solute species and mobile dislocations [8, 17] during deformation, and thus it constrains the dislocation motion to the slip plane. Consequently, in order to maintain an imposed strain rate, cyclic flow stress increases either to unlock the dislocations from the obstacles or to generate new dislocations. The moving dislocations get momentarily locked around the elastic strain fields of solute particles present in the matrix [1, 8] causing additional hardening of the matrix. At temperatures, where the effects of time dependent processes such as creep and oxidation are found to be minimal, the drastic reduction in LCF life has been reported to occur due to the deleterious effects of DSA in other alloys as well such as Nimonic PE 16 superalloy [9], Haynes 188 superalloy [10], Hastealloy [18] and Duplex stainless steel [19], 316 L(N) [20]. The cyclic hardening in the initial cycles observed in this alloy has been ascribed to the dislocation multiplication and accumulation within the slip band [21]. At elevated temperatures in addition to the dislocation multiplication and accumulation within the slip band [21] the continuous hardening is due to the formation of dislocation substructure which provides sites for numerous heterogeneous precipitations of fine M 23C6 particles. This in turn stabilizes the substructure during further cycling and prevents the recovery processes [22]. Srinivasan et al. [23] have observed an increase in friction stress with increasing temperature in the DSA regime and attributed to the interaction of solute atoms with mobile dislocations. The mechanism of the one/two/three slope behavior of CSR can be explained based on balance of factors affecting the deformation behavior. At ambient temperature, only dislocation multiplication occurs, whereas at elevated temperatures, the availability or non-availability of solutes for the interactions of solutes and dislocations decides the overall cyclic stress response. Detailed microstructural studies by Kaoumi et al. [24] recently confirmed that precipitate evolution occurs above 873 K, i.e. dissolution of the original primary carbides (MC or M6C formed during solidification of ingots) adds a significant amount of solute atoms into the matrix which act as barriers to dislocation motion, and leads to re-precipitation of smaller M23C6 carbides. At 1223 K, strengthening provided by dissolution of original primary carbides (MC or M6C) and reprecipitation of smaller M23C6 carbides competes with the softening process of recrystallization. The primary MC carbides present in the undissolved state during solution annealing have been reported to decompose into M23C6 and M6C on prolonged exposure at elevated temperatures [25]. At higher temperatures such as 973 K and above, a possibility of in-situ precipitation cannot be ruled out as has been reported to occur in Alloy 625 [16]. In Alloy 617, Wu et al. [26] indicated 973 K as the nose of the time-temperature-transformation diagram that favors formation of Ti(C, N), M23C23 and γ ׳phases. Maier et al. [27] recently reported that mere thermal aging of Alloy 617 at 973 K did not result in any significant microstructural changes, but under fatigue loading, significant amount of M 23C6 carbides precipitated within a relatively short time. The disappearance of serrations from hysteresis loops after few cycles at 973 K is associated with the unavailability of solute particles due to in-situ precipitate formation during the initial periods. Disappearance of serrations from flow curves have been also reported by Hayes [28] and Hayes and Hayes [29] in alloys like AISI 1020, 2.25 Cr-1Mo, alloy 718, and alloy 600 and by Vani Shankar et al. in Alloy 625 [16] during monotonic tensile tests. It has been suggested [30] that unlocking serrations are connected with precipitation before and during the test. Hayes [28] and Hayes and Hayes [29] carried out systematic investigations in AISI 1020, 2.25 Cr-1Mo, alloy 718, and alloy 600 to relate the strain for disappearance of serrations with strain rate and temperature. They have shown that the disappearance of serrations from the flow curve occurs in the high-temperature regime by either a progressively longer strain to the onset of serrations (a critical-strain-delay mechanism) or by a progressively smaller strain to the disappearance of serrations (disappearance off the end of the flow curve). In the former case, carbon diffusing down the dislocation line to a precipitate sink is suggested as the mechanism; however, in the latter case, carbon reacting with a carbide forming species on the dislocation line is responsible for the disappearance of 6
serrations. Thus, the important conclusion from their studies is that the disappearance is generally related to a precipitation mechanism, and the disappearance of serrated flow would occur when a balance is reached between the growth of carbon atmosphere and its depletion due to reaction between substitutional (carbideforming) atoms in the atmosphere. In this context, it has been suggested that the occurrence of type-C serrations could be regarded as a precursor to the precipitation and disappearance of serrations [30]. 4.2 Evolution of hysteresis loops and serrations: strength/magnitude of DSA interaction As stated earlier, the hysteresis loops showed serrations/stress drops in the plastic strained regions whose frequency of appearance and magnitude depended upon temperature-strain amplitude combination. Also, the serration heights varied during the life span of fatigue cycling. Since additional stresses are required to break free the solute-dislocation interaction atmospheres, the strength/magnitude of the interactions is evaluated by the magnitude of stress drops. The minimum strain required for first serration to appear in a hysteresis loop symbolizes how favorable it is to cause solute-dislocation interaction (DSA). Stress drop magnitude can be also expected to depend upon the solute species responsible for causing DSA. It is seen that the stress drops are significantly large during the initial few cycles and then the values saturates (Fig. 6 (a)). This denotes stronger interactions of solutes and dislocations during the initial periods of cyclic deformation and then formation of a stable microstructure/ solute-dislocation interaction zone. Once dislocations are generated during the initial deformation in a cycle, the diffusing solutes can interact with the moving dislocations and can cause strong solute-dislocation interactions. As the cyclic deformation progresses, the dislocations form a stable/constant microstructure in which the solutes come across almost constant number of dislocation atmospheres which results in stable stress or strain values after some time. Microstructural based quantification is however necessary to confirm the aforesaid. At 773 K, the magnitude of serrations is comparatively higher (among 573 and 773 K), and the minimum strain for first serration to appear is the lowest (among three temperatures, 573, 773 and 973 K). This accounts for the most favorable temperature for the occurrence of DSA at the strain rate of the LCF tests. The disappearance of serrations off the hysteresis loops after initial few cycles (Fig. 5 (c)) and delayed strains for the appearance of first serration at 973 K indicates that a different mechanism such as in-situ precipitation is operative. Numerous vacancies are generated during cyclic deformation that enhances in-situ precipitation. Another noteworthy feature is that there is a significant difference between the magnitude of serrations observed at 973 K than those at lower temperatures. Thus at 973 K the magnitude of serration is exceedingly higher (~ 125 MPa) than those observed at lower temperatures such as 773 K (10 to 60 MPa maximum). This amounts to the fact that the species responsible for causing DSA might be different at elevated temperatures such as 973 K than those operative at lower temperatures (below 973 K). Two temperature-strain rate combination regions in which two different species are responsible to cause DSA have been identified in a similar alloy system such as Alloy 625 under monotonic loading condition as well [16]. Thus whereas interstitial solutes such as carbon were responsible to cause DSA at lower temperatures, substitutional solute such as molybdenum caused DSA at higher temperatures in Alloy 625. It was stated that solutes were unavailable to cause further dislocation-solute interactions due to precipitation in Alloy 625 and hence serrations disappeared from the tensile flow curves after some time. Recently Ekaputra et al. [2] under monotonic loading condition has identified three zones for the occurrence of DSA in Alloy 617. Through activation energy calculation for solute diffusion in the nickel matrix, three temperature zones were identified for causing DSA. Thus Type D serrations occurred at 473 K (activation energy of 77.8 kJ/mol corresponded to pipe diffusion of carbon), Type A+B serrations between 573 - 773 K (activation energy ranged from 110.7 kJ/mol to 193.8 kJ/mol and was attributed to lattice diffusion of interstitial solute atoms of carbon) and Type C serrations that appeared between 873 to 973 K (activation energy of 265.1 kJ/mol, corresponded to the 7
lattice diffusion of substitutional solute atoms, such as chromium). In similar lines it may be stated that under cyclic loading conditions also, while carbon solutes causes DSA at lower temperatures, that at higher temperatures, Mo or Cr might be responsible. In-situ precipitation of M23C6 causes disappearance of serrations off the hysteresis loops after few cycles at 973 K. Detailed microscopic examination and determination of activation energy for the solutes responsible for causing DSA under cyclic loading is however necessary to confirm the aforesaid.
CONCLUSIONS 1) Occurrence of dynamic strain aging (DSA) in the temperature range 573 to 973 K is the cause for matrix hardening in Alloy 617M. 2) The magnitude and frequency of appearance of the serrations in hysteresis loops are linked with the strength of the solute-dislocation interactions causing DSA and is sensitive to strain amplitudetemperature combination. 3) 773 K is the most favored temperature for the occurrence of DSA in Alloy 617M in the range of temperatures and strain rate employed in this study. Two temperature regimes in which two species might be responsible to cause DSA in this alloy exists. At temperatures below 973 K, B type serrations are caused by interstitial carbon solutes whereas at 973 K other substitutional elements are responsible to cause C-type serrations and DSA. Appearance and then disappearance of the serrations after few cycles at 973 K indicate unavailability of the carbide forming element in the matrix possibly due to in-situ precipitation of M23C6.
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[17] A.W. Sleeswyk, Acta Metall6 (9)(1958)598–603. [18] M. G. Castelli, R. V. Miner and D. N. Robinson, ASTM-STP 1186, ASTM, Philadelphia, 1993;106 [19] S. Herenu, I. Alvarez-Armas and A. F. Armas, Scr. Mater. 45 (2001) 739-745. [20] V. S. Srinivasan, M. Valsan, R. Sandhya, K. Bhanu Sankara Rao, S. L. Mannan and D. H. Sastry, Int. J. Fatigue 21 (1999) 11-21. [21] G. Chai, P. Liu and J. Frodigh, J. Mater Sci., 39 (2004) 2689-2697. [22] M. A. Burke, and C. G. Beck, Met Trans A 15A (1984) 661-670. [23] V.S. Srinivasan, R. Sandhya, M. Valsan, K. Bhanu Sankara Rao, S.L. Mannan and D.H. Sastry , Scripta Materialia, 37(10)(1997) 1593-l 598. [24] D. Kaoumi, K. Hrutkay, Journal of Nuclear Materials 454 (2014) 265–273. [25] D.R. Muzyka: in The Superalloys 718, C.T. Sims and W.C. Hagel, eds., Wiley, New York, NY, 1972, p. 113. [26] Q. Wu, H. Song, R. W. Swindeman, Metallurgical and Materials Transactions 39A(11) (2008) 25692585. [27] Gerhard Maier, Hermann Riedel, Christoph Somsen, International Journal of Fatigue 55 (2013) 126–135. [28]R.W. Hayes, Acta Metall., 31 (1983) 365-371. [29] R.W. Hayes and W.C. Hayes: Acta Metall., 32(1984)259-267. [30] S. Venkadesan, C. Phaniraj, P.V. Sivaprasad, and P. Rodriguez, Acta Metall., 40 (1992)569-580.
Figure 1 Optical micrograph depicting initial microstructure of untested sample
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(a)
(b) Figure 2 Effect of strain amplitude on the cyclic stress response of Alloy 617M at (a) 300 K and 773 K.
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Figure 3 Effect of temperature on the cyclic stress response of Alloy 617M at +0.6% strain amplitude.
(a)
(b)
(c)
(d)
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(e) Figure 4 Effect of change in strain amplitude (at constant temperature such as at (a) 300 K, (b) 573 K, (c) 773 K and (d) 973 K) and (e) temperature (at constant strain amplitude, +0.4%) on the cyclic hardening evolution curve of Alloy 617 M.
(a)
(b)
(c) Figure 5 Representation of the evolution of hysteresis loops showing effects of temperature- through tests carried out at (a) 300 K; +0.6%, (b) 773 K; +0.6% and (c) 973 K; +0.6%. 12
(a)
(b)
(c) (d) Fig. 6 Plot of maximum stress drop in a cycle ((a) and (c)) and critical strain for appearance of first serration in hysteresis loop ((b) and (d)) versus the number of cycles showing for various strain amplitudes ((a) and (b)) and temperatures ((c) and (d)).
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Table 1 Chemical composition of Alloy 617M forging Element
C
Mo
Fe
Co
Ti
Cr
Si
Al
B
Ni
Wt% (F)
0.07
9.38
0.14
12.02
0.42
22.12
0.02
1.12
0.004
Bal.
Table 2 Summary of shape of cyclic stress response curve as a function of temperature-strain amplitude combination of the test
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Highlights 1) Dynamic strain aging (DSA) occurs between 573 and 973 K and causes matrix hardening. 2) Magnitude and frequency of appearance of serrations in hysteresis loops linked with strength of solutedislocation interactions. 3) 773 K is the most favored temperature for occurrence of DSA. 4) More than one species responsible for causing DSA.
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