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Influence of dwell times on the thermomechanical fatigue behavior of a directionally solidified Ni-base superalloy
International Journal of Fatigue 80 (2015) 426–433
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International Journal of Fatigue journal homepage: www.elsevier.com/locate/ijfatigue
Influence of dwell times on the thermomechanical fatigue behavior of a directionally solidified Ni-base superalloy Stefan Guth a,⇑, Roman Petráš b, Viktor Škorík b, Tomáš Kruml b, Jirˇí Man b, Karl-Heinz Lang a, Jaroslav Polák b a b
Institute for Applied Materials, Karlsruhe Institute of Technology, Engelbert-Arnold-Strasse 4, 76131 Karlsruhe, Germany Institute of Physics of Materials ASCR, Zˇizˇkova 22, 61662 Brno, Czech Republic
a r t i c l e
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Article history: Received 16 March 2015 Received in revised form 2 July 2015 Accepted 3 July 2015 Available online 8 July 2015 Keywords: Nickel base superalloy Thermomechanical fatigue Dwell time Lifetime behavior Damage mechanisms
a b s t r a c t In-phase (IP) and out-of-phase (OP) thermomechanical fatigue tests with T = 100–750 °C and optional dwells of 20 min at 750 °C were carried out on directionally solidified Ni-base Alloy 247 LC DS. Introducing dwells reduced the lifetime for both phase angles to about one sixth. Specific damage mechanisms were internal carbide and carbide–matrix interface cracks in IP tests and crack propagation along {1 1 1}-microtwin planes in OP tests. Introducing dwells intensified both effects, thus contributing to the lifetime reduction. During dwells, the gauge length may exhibit transversal creep because of extensometer forces distorting the strain measurement. Ó 2015 Elsevier Ltd. All rights reserved.
1. Introduction During service of gas turbine engines, the turbine blades are subject to various loading conditions. Temperature gradients during repeated start-up and shut down operations may induce thermomechanical fatigue (TMF) damage. Depending on the phase angle between mechanical loading and temperature, the total TMF damage consists of varying portions of fatigue, creep and environmental damage [1–3]. During steady state service, the exposure to high temperatures under load may degrade the material due to creep, oxidation and microstructural changes [4–9]. The interaction of temperature changes and applied cyclic strain during TMF and the degradation processes during steady state service often determines the service lifetime of the blades. In order to provide reliable lifetime predictions, a detailed knowledge about the material behavior under service conditions is necessary. In land-based gas turbines, the blades of the first and second stage are commonly made of directionally solidified (DS) Ni-base superalloys that contain a minimized amount of grain boundaries in the blade axis direction. That way, the creep resistance in this direction is improved and intergranular creep cavitation as well as preferential grain boundary oxidation is mostly suppressed. However, transgranular creep damage in a DS Ni-base superalloy was reported even at the relatively low temperature of 760 °C [10]. ⇑ Corresponding author. Tel.: +49 721 608 47450. E-mail address: [email protected] (S. Guth). http://dx.doi.org/10.1016/j.ijfatigue.2015.07.005 0142-1123/Ó 2015 Elsevier Ltd. All rights reserved.
c0 -strengthened DS and single crystal Ni-base alloys may exhibit anisotropic deformation behavior dependent on their orientation, temperature and strain rate [11–13]. Reasons for this are the orientation and temperature dependent tension–compression asymmetry of the c0 -phase [14,15] as well as time and temperature dependent dislocation reactions at c–c0 -interfaces [13]. More recently, microtwin deformation at intermediate temperatures was observed which may also result in anisotropic deformation behavior [16]. Dwell times representing steady state service typically reduce the lifetime of Ni-base superalloys in TMF or isothermal fatigue tests [17–19]; however positive effects have also been reported [20,21]. Because of the high costs of DS Ni-base alloys, laboratory tests are rare and only few literature data about their behavior under near service conditions are available. The objective of this study is to provide a better understanding about how dwell times influence the deformation, damage and lifetime behavior of Alloy 247 LC DS under TMF loading. Tests were conducted in laboratories of Karlsruhe Institute of Technology (KIT) and of Institute of Physics of Materials (IPM) in Brno. 2. Experimental details 2.1. Material The investigated material was Alloy 247 LC DS, a directionally solidified Ni-base superalloy which is primarily strengthened by c0 -precipitations. The alloy is comparable to CM247 LC DS that
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was studied earlier [10,22]. It is typically used for first and second stage turbine blades in land based gas turbines. The material was supplied in form of a casting slab from which solid round specimens were machined. The principal orientation of the directionally solidified grains was parallel to the specimen axis, i.e. the loading direction corresponds approximately to the [0 0 1] orientation. The elastic modulus at room temperature for this orientation is about 125 GPa. At KIT, specimens with a cylindrical gauge length of 30 mm and a gauge length diameter of 9 mm were tested. At IPM, a reduced gauge length of 14 mm and a gauge length diameter of 7 mm were used. The diameter of the columnar grains is rather heterogeneous ranging from a few hundred micrometers up to 2 mm, giving a relatively small number of grains within the gauge section. To allow for surface damage observation, the gauge sections of all specimens were polished. Initial polishing was done mechanically. It was followed by electrolytic (IPM) or oxide suspension (KIT) polishing methods. In Fig. 1, a SEM image of an electrolytically polished specimen surface is shown. The microstructure is well visible. The c0 -phase shows a bimodal distribution with fine dispersed precipitations in the matrix channels. Carbides as in the left lower corner could be frequently observed, particularly along the longitudinal grain boundaries.
(OPC), in-phase with dwell time (IPTD) and out-of-phase with dwell time (OPTD). In IP tests, the maximum temperature coincides with maximum mechanical strain, in OP tests with minimum mechanical strain. Accordingly, the dwell times occur under tensile load in IP tests and under compressive load in OP tests. During temperature transients, the temperature rate was 5 K/s giving a cycle time of 260 s in continuous tests and of 1460 s in tests with dwells. Prior to each test, the thermal strain eth during a cycle was derived by cycling the temperature at zero stress. The total strain– time waveform was calculated by adding the desired mechanical strain–time waveform to the thermal strain–time waveform (et = eth + emech). For comparability, the mechanical straining was symmetrical for both IP and OP cycling (Re = 1). The mechanical strain range was Demech = 1.5% for all tests resulting in a strain rate of e_ = 1.15 104 s1 during continuous cycling. The relatively high strain range was chosen to allow for macroscopic plastic strain and to yield test times within reasonable length. The lifetime was determined using a 10% drop of the stabilized maximum stress as failure criterion. To observe the development of internal damage and microstructure, some tests were stopped after a given number of cycles. Additionally, some tests were interrupted and restarted several times to study the surface evolution.
2.2. Thermomechanical fatigue testing
2.3. Microscopy
The experimental details in both laboratories were chosen as similar as possible. The TMF experiments were conducted on servo-hydraulic fatigue testing machines with 100 kN (KIT) or 50 kN (IPM) load cells. Induction systems were used for heating. Cooling was achieved by thermal conduction into the water cooled grips and could be additionally forced by air jets. For temperature measurement, ribbon type Ni–CrNi thermocouples (type K) were applied at the center of the gauge length. Total strain was measured with high temperature extensometers that were attached to the specimen using alumina tips. To avoid contact loss of the extensometer, the tips were borne against the specimen by a suspension system. The gauge length of the extensometer was 16 mm (KIT) and 12 mm (IPM). The tests were carried out under total strain control following the European Code of Practice [23]. The temperature range was 100–750 °C for all tests. 750 °C represents a moderate maximum service temperature experienced by blade materials in land based gas turbines. The minimum temperature of 100 °C is to represent a long shutdown of the turbine engine. Temperature and mechanical strain cycle had a triangular shape in continuous cycling and a trapezoidal shape in cycling with dwell times. The dwell time was 20 min at T = 750 °C. Phase angles between mechanical strain and temperature were 0° (in-phase, IP) and 180° (out-of-phase, OP). Thus, four test types were conducted: continuous in phase (IPC), continuous out-of-phase
TESCAN Lyra3 XMU FESEM with focused ion beam (FIB) and TESCAN Mira3 XM FESEM scanning electron microscopes (SEM) were used to observe surface damage and longitudinal sections. Microstructural changes were studied with a Philips CM12 transmission electron microscope (TEM). Thin foils were prepared using standard polishing procedures and were oriented using electron diffraction and Kikuchi lines.
Fig. 1. SEM image of the surface of an electropolished specimen prior to testing.
3. Results 3.1. Cyclic deformation and lifetime behavior Fig. 2 shows hysteresis loops during IPC and OPC cycling. In spite of the relatively high strain range the deformation is close to elastic. Since elastic modulus changes appreciably with temperature, the plastic strain amplitude was evaluated as the half-width of the hysteresis loop at mean stress. There is a negative mean stress in IPC cycles and a positive mean stress in OPC cycles. For a given cycle number, the absolute value of the mean stress is comparable for both phase angles. This suggests that for the tested strain rate there is no pronounced tension compression asymmetry of the cyclic yield stress at T = 100 and 750 °C. Therefore, the mean stress arises mainly due to the temperature dependence of the elastic modulus. For both phase angles, the plastic strain amplitude decreases slightly with increasing cycle number indicating cyclic hardening. Fig. 3 depicts hysteresis loops during IPTD and OPTD cycling. Compared to IPC and OPC tests, the plastic strain amplitude is higher due to stress relaxation during the dwells. Similar to the tests without dwells, there is a negative mean stress in the IPTD cycle and a positive mean stress in the OPTD cycle. The absolute value of the mean stress under OPTD loading was typically about twice the value under IPTD loading suggesting that plastic deformation at T = 750 °C is easier under compressive than under tensile stress. Since the absolute mean stress is comparable for IPC and OPC loading this anisotropic behavior is apparently induced by the dwells. During IPTD cycling, both the amount of relaxed stress during the dwells and the plastic strain amplitude decrease with increasing cycle number indicating cyclic hardening similar to IPC cycling. Contrarily, under OPTD loading, the relaxed stress
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Fig. 2. Stress–strain relationship for IP and OP cycling without dwell times.
Fig. 3. Stress–strain relationship for IPTD and OPTD cycling.
during dwells and the plastic strain amplitude increase with the cycle number. Fig. 4 shows the development of the absolute stress during the dwells of an OPTD test in the first cycle and at half of the lifetime. The exponential stress drop at half of the lifetime compared to the nearly linear decrease in the first cycle indicates that the operative relaxation mechanisms change during cycling. Although the increasing plastic strain amplitude under OPTD loading implies that the material cyclically softens, the stress amplitude increases simultaneously. This is because the slopes of the elastic parts of the
Fig. 4. Stress relaxation during OPTD dwell times.
OPTD loop increase during cycling. This effect was observed in both laboratories. After 250 OPTD cycles, a specimen showed an 18 % increase of stiffness (elastic modulus) at room temperature when measured with load cell and extensometer. Gauging of an OPTD tested specimen showed that the entire gauge length crept horizontally about 0.1 mm toward the direction in which the extensometer tips point. The corresponding process is illustrated in Fig. 5. The force applied by the extensometer suspension system is small but high enough to induce a transverse repeated creep
Fig. 5. Schematic of the transverse creep deformation of the specimen gauge length during dwell times in OPTD tests. The compressive forces in loading direction facilitate bending due to forces exerted by the extensometer.
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deformation during dwell times at 750 °C. The gauge length of the specimen bends gradually in the direction of applied force. During unloading from tension and loading into compression a fraction of the applied strain is due to bending and the measured strain range is smaller. Thus, the measured specimen stiffness increases. The effect occurred only during OPTD cycling when compressive stress at high temperature promotes bending during repeated plastic strain relaxation. In IP tests, tensile stress at high temperatures counteracts the bending while in OPC tests, the transverse creep during continuous loading is negligible. Inaccurate strain measurement due to specimen bending during dwells may occur for any similar testing situation in which an extensometer with suspension system is used and compressive
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dwells at elevated temperatures are applied. To avoid this, the use of self-supporting extensometers which apply no net force to the specimen is suggested. The lifetimes for the different test conditions are compared in Fig. 6. For tests without dwells, IP lifetime is about one third of OP lifetime. Introducing dwells reduced both IP and OP lifetime to about one sixth. The lifetimes of tests with dwells measured in both laboratories are very close. It should be noted that the tests at IPM were interrupted several times. 3.2. Microstructure Fig. 7 shows typical TEM micrographs of specimens after OPC and OPTD cycling. For both cycle types, there is a high dislocation density in the matrix channels. In many cases superdislocations belonging to {1 1 1}h1 1 0i slip systems sheared c0 -precipitates. Crystallographic bands along the traces of {1 1 1}-planes cutting both matrix and c0 -phase could be frequently observed. The bands are more numerous and distinct after OPTD cycling as shown in Fig. 7b. The diffraction pattern from an area close to such lines in Fig. 7c reveals twin reflections proving that the deformation bands correspond to microtwins. In Fig. 8a and b, typical microstructures after IPC and IPTD tests are presented. The dislocation structures in the matrix and within the c0 -precipitates are similar to those in OPC and OPTD tested specimens (see Fig. 7). Microtwins were not found. Rafting of c0 -precipitates did not occur at all conditions tested. 3.3. Surface and internal damage
Fig. 6. Lifetime of specimens cycled with constant strain range Demech = 1.5% under different regimes. The tests at IPM were interrupted several times.
Fig. 9 shows representative SEM images of surface damage after OP cycling. Crack densities of OP tested specimens were generally
Fig. 7. Representative TEM images of OP cycled specimens revealing microtwins: (a) after OPC cycling, (b) after OPTD cycling, and (c) diffraction pattern in the area of a microtwin in image (b), the zone axis is [1 1 0]; twin reflections are denoted by a T.
Fig. 8. Representative TEM images of IP cycled specimens: (a) after IPC cycling, and (b) after IPTD cycling.
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Fig. 9. Surface SEM images of OP cycled specimens: (a) secondary cracks after OPC cycling, (b) crack initiation at coarse carbides after OPTD cycling, and (c) deformation lines after OPTD cycling. The loading axis is horizontal.
Fig. 10. Cracked specimen after OPTD cycling. The main crack propagated partly under about 45° to the horizontal loading axis.
high and the crack morphology was rather flat. Fig. 9a shows secondary surface cracks on a specimen after OPC cycling. The cracks are seamed with oxides and propagate mainly perpendicularly to the loading axis. Some cracks (e.g. the longest crack in Fig. 9a) change their direction to about 45° to the loading axis. In OPTD tests, the surface carbides coarsened significantly. As can be seen
in Fig. 9b and c, cracks initiate typically at coarsened carbides and propagate then into the bulk material. Many of the coarsened carbides were already cracked after less than 20% of the lifetime. EDX analysis revealed that the carbides contain a high amount of Hf and Ta. During OPTD tests, distinct deformation lines with an inclination angle of approximately 45° to the loading axis developed in some areas of the gauge section as shown in Fig. 9c. Although the deformation lines created a surface relief with notch-like intrusions, most secondary cracks initiated and propagated under an angle of about 90° to the loading axis. However, as shown in Fig. 10, the main cracks of specimens cycled until failure propagated in large part under an angle to the loading axis that is similar to the inclination angle of the deformation lines. After OP cycling, no internal damage could be found. The number of surface cracks after IP loading was generally much lower than after OP loading. Particularly after IPC cycling very few secondary surface cracks could be found. However, IP cracks were typically longer than OP cracks. Fig. 11a shows surface cracks on an IPTD cycled specimen. The morphology of the cracks is wavier when compared to typical OP cracks. In IPTD tests, surface carbides coarsened similar to OPTD tests but were typically not cracked. After both IPC and IPTD tests, internal damage in form of cracked carbides and carbide–matrix interfaces could be frequently found. Fig. 11b shows representative internal damage on a longitudinal section of a specimen after IPTD cycling. Apparently, the cracks tend to coalesce. Introducing dwell times significantly increased the extent of internal damage under IP loading.
Fig. 11. SEM images of specimens after IP cycling with dwell times: (a) surface cracks, and (b) electropolished longitudinal section showing internal damage. The loading axis is horizontal.
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4. Discussion 4.1. Microstructure and deformation behavior The TEM investigations revealed that depending on the TMF phase angle, three mechanisms of dislocation movement are possible. Ordinary dislocation movement in the c-channels and shearing of c0 -precipitates by dislocations in {1 1 1}h1 1 0i slip system occur similarly during IP and OP cycling. Introducing dwell times has no significant influence on these mechanisms. Microtwin formation occurs only under OP loading and is intensified by dwell times. Microtwinning in Ni-base superalloys has been reported as an important deformation mechanism in the temperature range between 650 and 850 °C [24,25]. Recently, it has been observed after TMF OP tests of CMSX-4 single crystals [26–28]. Kakehi [16] pointed out that twin deformation in Ni-base single crystals loaded along [0 0 1] orientation occurs preferentially under compression since twinning shear on {1 1 1} planes in h1 1 2i direction is favored by compressive stress. The model proposed by Kolbe [24] assumes that the microtwins are formed by the passage of pairwise a/6 h1 1 2i Shockley partials on adjacent {1 1 1} planes. Experimental results of Kovarik et al. [25] support this assumption. In the ordered L12 structure of the c0 -phase having nominal stoichiometry Ni3Al, shearing of two a/6 h1 1 2i Shockley partials creates an energetically unfavorable pseudotwin with Al atoms as nearest neighbors. These pseudotwins should be converted into a true twin by a diffusion mediated reordering process. The reordering process is seen as rate-limiting for the microtwin deformation [25]. The temperature dependence of diffusion explains why microtwins form only at elevated temperatures. Thus, microtwins can form when compressive stress coincides with elevated temperatures, which is given in the case of OP loading. Since dwells allow for intensified diffusion reordering, their introduction will increase the extent of microtwinning as observed after OPTD tests. Thus, the present results are in agreement with the model of Kolbe [24] and the assumption that the diffusion reordering is rate limiting for the microtwin deformation. It has been reported that microtwinning results in easier creep deformation [16] and a reduced critical resolved shear stress [24]. From this can be inferred that the observed easier plastic deformation under compressive load at 750 °C resulting in a higher absolute mean stress in OPTD tests than in IPTD tests is produced by extensive microtwin formation. The cyclic increase of the relaxed stress along with the apparent change of relaxation mechanism during OPTD dwells (Fig. 4) may also be attributed to microtwin deformation. Localized plastic deformation along the {1 1 1} microtwin planes produces the surface deformation lines observed after OPTD loading whose orientation corresponds approximately to the trace of the {1 1 1} twin planes (Fig. 9c). Based on TMF OP experiments with dwell times on Ni-base single crystals, this was also concluded by Zhang et al. [29]. It is noticeable that during OPTD loading, the tensile cyclic yield stress decreased with increasing cycle number although under tensile loading at temperatures around 100 °C, microtwinning should not occur (Fig. 3). Microtwins along {1 1 1} planes apparently facilitate ordinary dislocation movement in {1 1 1}h1 1 0i slip system. Assuming that the loading axis is parallel to the initial [0 0 1] direction and the 1 1Þ, the twinned lattice would be loaded twin mirror plane is ð1 1 direction. The maximum Schmid factor for along the ½2 2 and ½2 2 1 is 0.408. {1 1 1}h1 1 0i slip for both loading along ½0 0 1 The Schmid factors of the cubic {1 0 0}h1 1 0i systems along ½0 0 1 1 are generally higher. According to is 0 and the values along ½2 2 the model of Lall et al. [15], this would result in an increased critical resolved shear stress in the twinned c0 lattice which is contrary to observations. Consequently, with this idealized consideration,
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the cyclic softening cannot be attributed to a change of crystal orientation due to the twins. Possibly, the necessary stress for dislocation slip across c–c0 interfaces is reduced when the interface is twinned. However, this needs to be investigated in further work. The cyclic hardening during IPC, OPC and particularly IPTD cycling (Figs. 2 and 3) can be explained by dislocation multiplication in the c-channels impeding further dislocation movement. The dislocation structure in the c-channels indicates that the same hardening effect takes also place in OPTD tests. However, in this case it is outweighed by softening due to microtwinning. In OPC tests, the amount of microtwins formed is apparently too low to cause macroscopic softening. 4.2. Damage mechanisms and lifetime behavior Although constant mechanical strain range has been applied for the four test types, the lifetimes differ significantly. The reason for this is that the dominant damage mechanisms vary with the TMF phase angle and may be enhanced by dwell times. Furthermore, additional damage mechanisms may operate during the dwells. In OP tests, the tensile half-cycle occurs at low temperature when few slip systems are active and the surface oxide layer has low ductility. Thus, cracks initiate easily at the surface oxides and propagate in low plasticity regime resulting in a flat and sharp morphology as shown in Fig. 9a. It has been frequently stated that deformation microtwins may provide a preferential crack path along which crack propagation accelerates [26–30]. Hence, principal cracks under approximately 45° to the loading axis and the transition of crack propagation from perpendicular to approximately 45° to the loading axis may be attributed to crack growth along the {1 1 1} microtwin planes. The introduction of dwell times leads to more frequent microtwinning which is accompanied by material softening. Because of the resulting higher plastic strain amplitude and more possible crack paths along microtwin planes, cracks grow more rapidly contributing to the observed decrease of fatigue lifetime due to dwells. Zhang et al. [30] found that the portion of crystallographic cracking along microtwins increases with the introduction of dwells in TMF-OP tests which supports this assumption. Since compressive stress acts during OPTD dwells, the cracks are not opened and oxidation of the crack tip area should have a minor effect on OPTD lifetime. However, surface oxidation and coarsening of superficial carbides during the dwells may promote the formation of preferential crack initiation sites. That way, the stage of crack initiation is shortened. Under IP loading, the tensile half-cycle occurs at elevated temperatures when the surface oxide layer has higher ductility. Therefore, the surface oxide layer has only minimum effect on crack initiation [22]. This is supported by the results of Gordon et al. [4] stating that for a given mechanical strain range, the critical oxide rupture thickness is significantly higher for IP than for OP loading. Similarly, coarse surface carbides developing in IPTD tests play a minor role for crack initiation. Consequently, IP specimens exhibit fewer surface cracks than OP specimens. Since at elevated temperatures more slip systems are active, IP cracks can propagate by repeated plastic blunting and re-sharpening which is the reason for their wavy morphology (Fig. 11a) [27]. The lifetime reduction due to dwells under IP loading may be attributed to various mechanisms. The increased internal damage after IPTD tests indicates that surface cracks linking up with internal cracks may accelerate their propagation rate. Creep crack propagation results of a Ni-base DS alloy at 760 °C suggest that cracks may grow substantially during the tensile dwells [10]. During IP dwells, the tensile loading opens up the cracks and the crack tips are exposed to ambient oxygen. Thus, oxidation induced formation of an embrittled zone [6,7] and c0 -dissolving [31] ahead of the crack tip during dwells may facilitate crack propagation in the following cycles.
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For both tests with and without dwells, IP lifetime was shorter than OP lifetime although OP loading results in a higher maximum stress. In case of OPTD tests, the plastic strain amplitude is also higher. The results are in agreement with reported shorter IP lives at high mechanical strain ranges for other Ni-base superalloys [2,32–34]. Considering that cracks initiate early in OP lifetime, it can be inferred that the average crack propagation rate under IP conditions is significantly higher than under comparable OP conditions. The observed numerous short surface cracks after OP loading compared to few long cracks after IP loading support this assumption. Assuming that cracks propagate predominantly in the tensile part of the cycle, this suggests that crack growth at high temperatures is much more rapid than at low temperatures. This may be related to the more homogeneous dislocation slip at high temperatures compared to planar slip at low temperatures [35]. He et al. [36] recently reported an increase of the growth rate of small cracks with temperature for a Ni-base DS alloy. They attributed this to material softening at elevated temperatures. At high temperatures, cracks are opened under IP loading and closed under OP loading. Hence, the role of oxidation induced embrittlement and c0 -dissolving ahead of the crack tip is presumably much more pronounced under IP loading which may be a further reason for higher IP crack propagation rates. Because of the complex interaction of fatigue, creep and oxidation damage combined with twin formation and stress redistributions, it is difficult to assess which damage mechanism governs the lifetime for the four test types, respectively. In order to quantify the contributions of the particular damage mechanisms, further work including tests with varying dwell time lengths and mechanical strain amplitudes is needed.
5. Conclusions The effects of 20 min dwell times on the deformation, damage and lifetime behavior of the directionally solidified Ni-base alloy Alloy 247 LC DS in IP and OP TMF tests with constant strain range of 1.5% and a temperature range of 100–750 °C has been studied. The following conclusions can be drawn: 1. Lifetime under IP loading is about one third of the lifetime under OP loading. For both phase angles, the introduction of 20 min dwells at 750 °C reduced the lifetime to about one sixth. 2. During OP cycling, microtwin deformation occurs which is intensified by dwell times. Microtwins apparently facilitate plastic deformation along {1 1 1} planes resulting in material softening. After IP loading, no microtwinning was observed. The crystallography/diffusion-based model for microtwinning introduced by Kolbe [24] could explain the dependence on phase angle and dwell times. 3. Fatigue crack propagation along microtwins in OP tests and formation of internal cracks at carbides and carbide/matrix interfaces in IP tests represent specific damage mechanisms identified in TMF of a directionally solidified Ni-base superalloy. Introducing dwells intensifies both mechanisms thus contributing to the observed lifetime reduction. Increased oxidation during dwells promoting superficial crack initiation under OP loading and embrittlement as well as c0 -dissolving ahead of the crack tip under IP loading may further reduce the lifetime. 4. During compressive dwell times of OP tests, the gauge length may exhibit transverse creep deformation under action of the force applied by the extensometer rods. That way the strain measurement is distorted which leads to an apparent increase of specimen stiffness. In order to avoid this, the use of self-supporting extensometers is suggested for any similar testing situation including compressive dwells at elevated temperatures.
Acknowledgements This work was supported by the Ministry of Education, Youth and Sports of the Czech Republic throughout the project No. CZ.1.07/2.3.00/30.0063 ‘‘Talented postdocs for scientific excellence in physics of materials‘‘. The financial support of the Karlsruhe House of Young Scientists (KHYS) for the stay abroad at the Institute of Physics of Materials (IPM) in Brno is gratefully acknowledged. Sincere thanks to Ludvik Kunz for the invitation to the research stay.
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S. Guth et al. / International Journal of Fatigue 80 (2015) 426–433 [25] Kovarik L, Unocic RR, Li Ju, Sarosi P, Shen C, Wang Y, Mills MJ. Microtwinning and other shearing mechanisms at intermediate temperatures in Ni-based superalloys. Prog Mater Sci 2009;54:839–73. [26] Moverare JJ, Johansson S, Reed RC. Deformation and damage mechanisms during thermal–mechanical fatigue of a single-crystal superalloy. Acta Mater 2009;57:2266–76. [27] Hong HU, Kang JG, Choi BG, Kim IS, Yoo YS, Jo CY. A comparative study on thermomechanical and low cycle fatigue failures of a single crystal nickelbased superalloy. Int J Fatigue 2011;33:1592–9. [28] Hong HU, Yoon JG, Choi BG, Kim IS, Yoo YS, Jo CY. Localized microtwin formation and failure during out-of-phase thermomechanical fatigue of a single crystal nickel-based superalloy. Int J Fatigue 2014;69:22–7. [29] Zhang JX, Harada H, Ro Y, Koizumi Y, Kobayashi T. Thermomechanical fatigue mechanism in a modern single crystal nickel base superalloy TMS-82. Acta Mater 2008;56:2975–87. [30] Zhang JX, Harada H, Koizumi Y, Kobayashi T. Crack appearance of singlecrystal nickel-base superalloys after thermomechanical fatigue. Scripta Mater 2009;61:1105–8.
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Contents lists available at ScienceDirect
International Journal of Fatigue journal homepage: www.elsevier.com/locate/ijfatigue
Influence of dwell times on the thermomechanical fatigue behavior of a directionally solidified Ni-base superalloy Stefan Guth a,⇑, Roman Petráš b, Viktor Škorík b, Tomáš Kruml b, Jirˇí Man b, Karl-Heinz Lang a, Jaroslav Polák b a b
Institute for Applied Materials, Karlsruhe Institute of Technology, Engelbert-Arnold-Strasse 4, 76131 Karlsruhe, Germany Institute of Physics of Materials ASCR, Zˇizˇkova 22, 61662 Brno, Czech Republic
a r t i c l e
i n f o
Article history: Received 16 March 2015 Received in revised form 2 July 2015 Accepted 3 July 2015 Available online 8 July 2015 Keywords: Nickel base superalloy Thermomechanical fatigue Dwell time Lifetime behavior Damage mechanisms
a b s t r a c t In-phase (IP) and out-of-phase (OP) thermomechanical fatigue tests with T = 100–750 °C and optional dwells of 20 min at 750 °C were carried out on directionally solidified Ni-base Alloy 247 LC DS. Introducing dwells reduced the lifetime for both phase angles to about one sixth. Specific damage mechanisms were internal carbide and carbide–matrix interface cracks in IP tests and crack propagation along {1 1 1}-microtwin planes in OP tests. Introducing dwells intensified both effects, thus contributing to the lifetime reduction. During dwells, the gauge length may exhibit transversal creep because of extensometer forces distorting the strain measurement. Ó 2015 Elsevier Ltd. All rights reserved.
1. Introduction During service of gas turbine engines, the turbine blades are subject to various loading conditions. Temperature gradients during repeated start-up and shut down operations may induce thermomechanical fatigue (TMF) damage. Depending on the phase angle between mechanical loading and temperature, the total TMF damage consists of varying portions of fatigue, creep and environmental damage [1–3]. During steady state service, the exposure to high temperatures under load may degrade the material due to creep, oxidation and microstructural changes [4–9]. The interaction of temperature changes and applied cyclic strain during TMF and the degradation processes during steady state service often determines the service lifetime of the blades. In order to provide reliable lifetime predictions, a detailed knowledge about the material behavior under service conditions is necessary. In land-based gas turbines, the blades of the first and second stage are commonly made of directionally solidified (DS) Ni-base superalloys that contain a minimized amount of grain boundaries in the blade axis direction. That way, the creep resistance in this direction is improved and intergranular creep cavitation as well as preferential grain boundary oxidation is mostly suppressed. However, transgranular creep damage in a DS Ni-base superalloy was reported even at the relatively low temperature of 760 °C [10]. ⇑ Corresponding author. Tel.: +49 721 608 47450. E-mail address: [email protected] (S. Guth). http://dx.doi.org/10.1016/j.ijfatigue.2015.07.005 0142-1123/Ó 2015 Elsevier Ltd. All rights reserved.
c0 -strengthened DS and single crystal Ni-base alloys may exhibit anisotropic deformation behavior dependent on their orientation, temperature and strain rate [11–13]. Reasons for this are the orientation and temperature dependent tension–compression asymmetry of the c0 -phase [14,15] as well as time and temperature dependent dislocation reactions at c–c0 -interfaces [13]. More recently, microtwin deformation at intermediate temperatures was observed which may also result in anisotropic deformation behavior [16]. Dwell times representing steady state service typically reduce the lifetime of Ni-base superalloys in TMF or isothermal fatigue tests [17–19]; however positive effects have also been reported [20,21]. Because of the high costs of DS Ni-base alloys, laboratory tests are rare and only few literature data about their behavior under near service conditions are available. The objective of this study is to provide a better understanding about how dwell times influence the deformation, damage and lifetime behavior of Alloy 247 LC DS under TMF loading. Tests were conducted in laboratories of Karlsruhe Institute of Technology (KIT) and of Institute of Physics of Materials (IPM) in Brno. 2. Experimental details 2.1. Material The investigated material was Alloy 247 LC DS, a directionally solidified Ni-base superalloy which is primarily strengthened by c0 -precipitations. The alloy is comparable to CM247 LC DS that
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was studied earlier [10,22]. It is typically used for first and second stage turbine blades in land based gas turbines. The material was supplied in form of a casting slab from which solid round specimens were machined. The principal orientation of the directionally solidified grains was parallel to the specimen axis, i.e. the loading direction corresponds approximately to the [0 0 1] orientation. The elastic modulus at room temperature for this orientation is about 125 GPa. At KIT, specimens with a cylindrical gauge length of 30 mm and a gauge length diameter of 9 mm were tested. At IPM, a reduced gauge length of 14 mm and a gauge length diameter of 7 mm were used. The diameter of the columnar grains is rather heterogeneous ranging from a few hundred micrometers up to 2 mm, giving a relatively small number of grains within the gauge section. To allow for surface damage observation, the gauge sections of all specimens were polished. Initial polishing was done mechanically. It was followed by electrolytic (IPM) or oxide suspension (KIT) polishing methods. In Fig. 1, a SEM image of an electrolytically polished specimen surface is shown. The microstructure is well visible. The c0 -phase shows a bimodal distribution with fine dispersed precipitations in the matrix channels. Carbides as in the left lower corner could be frequently observed, particularly along the longitudinal grain boundaries.
(OPC), in-phase with dwell time (IPTD) and out-of-phase with dwell time (OPTD). In IP tests, the maximum temperature coincides with maximum mechanical strain, in OP tests with minimum mechanical strain. Accordingly, the dwell times occur under tensile load in IP tests and under compressive load in OP tests. During temperature transients, the temperature rate was 5 K/s giving a cycle time of 260 s in continuous tests and of 1460 s in tests with dwells. Prior to each test, the thermal strain eth during a cycle was derived by cycling the temperature at zero stress. The total strain– time waveform was calculated by adding the desired mechanical strain–time waveform to the thermal strain–time waveform (et = eth + emech). For comparability, the mechanical straining was symmetrical for both IP and OP cycling (Re = 1). The mechanical strain range was Demech = 1.5% for all tests resulting in a strain rate of e_ = 1.15 104 s1 during continuous cycling. The relatively high strain range was chosen to allow for macroscopic plastic strain and to yield test times within reasonable length. The lifetime was determined using a 10% drop of the stabilized maximum stress as failure criterion. To observe the development of internal damage and microstructure, some tests were stopped after a given number of cycles. Additionally, some tests were interrupted and restarted several times to study the surface evolution.
2.2. Thermomechanical fatigue testing
2.3. Microscopy
The experimental details in both laboratories were chosen as similar as possible. The TMF experiments were conducted on servo-hydraulic fatigue testing machines with 100 kN (KIT) or 50 kN (IPM) load cells. Induction systems were used for heating. Cooling was achieved by thermal conduction into the water cooled grips and could be additionally forced by air jets. For temperature measurement, ribbon type Ni–CrNi thermocouples (type K) were applied at the center of the gauge length. Total strain was measured with high temperature extensometers that were attached to the specimen using alumina tips. To avoid contact loss of the extensometer, the tips were borne against the specimen by a suspension system. The gauge length of the extensometer was 16 mm (KIT) and 12 mm (IPM). The tests were carried out under total strain control following the European Code of Practice [23]. The temperature range was 100–750 °C for all tests. 750 °C represents a moderate maximum service temperature experienced by blade materials in land based gas turbines. The minimum temperature of 100 °C is to represent a long shutdown of the turbine engine. Temperature and mechanical strain cycle had a triangular shape in continuous cycling and a trapezoidal shape in cycling with dwell times. The dwell time was 20 min at T = 750 °C. Phase angles between mechanical strain and temperature were 0° (in-phase, IP) and 180° (out-of-phase, OP). Thus, four test types were conducted: continuous in phase (IPC), continuous out-of-phase
TESCAN Lyra3 XMU FESEM with focused ion beam (FIB) and TESCAN Mira3 XM FESEM scanning electron microscopes (SEM) were used to observe surface damage and longitudinal sections. Microstructural changes were studied with a Philips CM12 transmission electron microscope (TEM). Thin foils were prepared using standard polishing procedures and were oriented using electron diffraction and Kikuchi lines.
Fig. 1. SEM image of the surface of an electropolished specimen prior to testing.
3. Results 3.1. Cyclic deformation and lifetime behavior Fig. 2 shows hysteresis loops during IPC and OPC cycling. In spite of the relatively high strain range the deformation is close to elastic. Since elastic modulus changes appreciably with temperature, the plastic strain amplitude was evaluated as the half-width of the hysteresis loop at mean stress. There is a negative mean stress in IPC cycles and a positive mean stress in OPC cycles. For a given cycle number, the absolute value of the mean stress is comparable for both phase angles. This suggests that for the tested strain rate there is no pronounced tension compression asymmetry of the cyclic yield stress at T = 100 and 750 °C. Therefore, the mean stress arises mainly due to the temperature dependence of the elastic modulus. For both phase angles, the plastic strain amplitude decreases slightly with increasing cycle number indicating cyclic hardening. Fig. 3 depicts hysteresis loops during IPTD and OPTD cycling. Compared to IPC and OPC tests, the plastic strain amplitude is higher due to stress relaxation during the dwells. Similar to the tests without dwells, there is a negative mean stress in the IPTD cycle and a positive mean stress in the OPTD cycle. The absolute value of the mean stress under OPTD loading was typically about twice the value under IPTD loading suggesting that plastic deformation at T = 750 °C is easier under compressive than under tensile stress. Since the absolute mean stress is comparable for IPC and OPC loading this anisotropic behavior is apparently induced by the dwells. During IPTD cycling, both the amount of relaxed stress during the dwells and the plastic strain amplitude decrease with increasing cycle number indicating cyclic hardening similar to IPC cycling. Contrarily, under OPTD loading, the relaxed stress
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Fig. 2. Stress–strain relationship for IP and OP cycling without dwell times.
Fig. 3. Stress–strain relationship for IPTD and OPTD cycling.
during dwells and the plastic strain amplitude increase with the cycle number. Fig. 4 shows the development of the absolute stress during the dwells of an OPTD test in the first cycle and at half of the lifetime. The exponential stress drop at half of the lifetime compared to the nearly linear decrease in the first cycle indicates that the operative relaxation mechanisms change during cycling. Although the increasing plastic strain amplitude under OPTD loading implies that the material cyclically softens, the stress amplitude increases simultaneously. This is because the slopes of the elastic parts of the
Fig. 4. Stress relaxation during OPTD dwell times.
OPTD loop increase during cycling. This effect was observed in both laboratories. After 250 OPTD cycles, a specimen showed an 18 % increase of stiffness (elastic modulus) at room temperature when measured with load cell and extensometer. Gauging of an OPTD tested specimen showed that the entire gauge length crept horizontally about 0.1 mm toward the direction in which the extensometer tips point. The corresponding process is illustrated in Fig. 5. The force applied by the extensometer suspension system is small but high enough to induce a transverse repeated creep
Fig. 5. Schematic of the transverse creep deformation of the specimen gauge length during dwell times in OPTD tests. The compressive forces in loading direction facilitate bending due to forces exerted by the extensometer.
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deformation during dwell times at 750 °C. The gauge length of the specimen bends gradually in the direction of applied force. During unloading from tension and loading into compression a fraction of the applied strain is due to bending and the measured strain range is smaller. Thus, the measured specimen stiffness increases. The effect occurred only during OPTD cycling when compressive stress at high temperature promotes bending during repeated plastic strain relaxation. In IP tests, tensile stress at high temperatures counteracts the bending while in OPC tests, the transverse creep during continuous loading is negligible. Inaccurate strain measurement due to specimen bending during dwells may occur for any similar testing situation in which an extensometer with suspension system is used and compressive
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dwells at elevated temperatures are applied. To avoid this, the use of self-supporting extensometers which apply no net force to the specimen is suggested. The lifetimes for the different test conditions are compared in Fig. 6. For tests without dwells, IP lifetime is about one third of OP lifetime. Introducing dwells reduced both IP and OP lifetime to about one sixth. The lifetimes of tests with dwells measured in both laboratories are very close. It should be noted that the tests at IPM were interrupted several times. 3.2. Microstructure Fig. 7 shows typical TEM micrographs of specimens after OPC and OPTD cycling. For both cycle types, there is a high dislocation density in the matrix channels. In many cases superdislocations belonging to {1 1 1}h1 1 0i slip systems sheared c0 -precipitates. Crystallographic bands along the traces of {1 1 1}-planes cutting both matrix and c0 -phase could be frequently observed. The bands are more numerous and distinct after OPTD cycling as shown in Fig. 7b. The diffraction pattern from an area close to such lines in Fig. 7c reveals twin reflections proving that the deformation bands correspond to microtwins. In Fig. 8a and b, typical microstructures after IPC and IPTD tests are presented. The dislocation structures in the matrix and within the c0 -precipitates are similar to those in OPC and OPTD tested specimens (see Fig. 7). Microtwins were not found. Rafting of c0 -precipitates did not occur at all conditions tested. 3.3. Surface and internal damage
Fig. 6. Lifetime of specimens cycled with constant strain range Demech = 1.5% under different regimes. The tests at IPM were interrupted several times.
Fig. 9 shows representative SEM images of surface damage after OP cycling. Crack densities of OP tested specimens were generally
Fig. 7. Representative TEM images of OP cycled specimens revealing microtwins: (a) after OPC cycling, (b) after OPTD cycling, and (c) diffraction pattern in the area of a microtwin in image (b), the zone axis is [1 1 0]; twin reflections are denoted by a T.
Fig. 8. Representative TEM images of IP cycled specimens: (a) after IPC cycling, and (b) after IPTD cycling.
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Fig. 9. Surface SEM images of OP cycled specimens: (a) secondary cracks after OPC cycling, (b) crack initiation at coarse carbides after OPTD cycling, and (c) deformation lines after OPTD cycling. The loading axis is horizontal.
Fig. 10. Cracked specimen after OPTD cycling. The main crack propagated partly under about 45° to the horizontal loading axis.
high and the crack morphology was rather flat. Fig. 9a shows secondary surface cracks on a specimen after OPC cycling. The cracks are seamed with oxides and propagate mainly perpendicularly to the loading axis. Some cracks (e.g. the longest crack in Fig. 9a) change their direction to about 45° to the loading axis. In OPTD tests, the surface carbides coarsened significantly. As can be seen
in Fig. 9b and c, cracks initiate typically at coarsened carbides and propagate then into the bulk material. Many of the coarsened carbides were already cracked after less than 20% of the lifetime. EDX analysis revealed that the carbides contain a high amount of Hf and Ta. During OPTD tests, distinct deformation lines with an inclination angle of approximately 45° to the loading axis developed in some areas of the gauge section as shown in Fig. 9c. Although the deformation lines created a surface relief with notch-like intrusions, most secondary cracks initiated and propagated under an angle of about 90° to the loading axis. However, as shown in Fig. 10, the main cracks of specimens cycled until failure propagated in large part under an angle to the loading axis that is similar to the inclination angle of the deformation lines. After OP cycling, no internal damage could be found. The number of surface cracks after IP loading was generally much lower than after OP loading. Particularly after IPC cycling very few secondary surface cracks could be found. However, IP cracks were typically longer than OP cracks. Fig. 11a shows surface cracks on an IPTD cycled specimen. The morphology of the cracks is wavier when compared to typical OP cracks. In IPTD tests, surface carbides coarsened similar to OPTD tests but were typically not cracked. After both IPC and IPTD tests, internal damage in form of cracked carbides and carbide–matrix interfaces could be frequently found. Fig. 11b shows representative internal damage on a longitudinal section of a specimen after IPTD cycling. Apparently, the cracks tend to coalesce. Introducing dwell times significantly increased the extent of internal damage under IP loading.
Fig. 11. SEM images of specimens after IP cycling with dwell times: (a) surface cracks, and (b) electropolished longitudinal section showing internal damage. The loading axis is horizontal.
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4. Discussion 4.1. Microstructure and deformation behavior The TEM investigations revealed that depending on the TMF phase angle, three mechanisms of dislocation movement are possible. Ordinary dislocation movement in the c-channels and shearing of c0 -precipitates by dislocations in {1 1 1}h1 1 0i slip system occur similarly during IP and OP cycling. Introducing dwell times has no significant influence on these mechanisms. Microtwin formation occurs only under OP loading and is intensified by dwell times. Microtwinning in Ni-base superalloys has been reported as an important deformation mechanism in the temperature range between 650 and 850 °C [24,25]. Recently, it has been observed after TMF OP tests of CMSX-4 single crystals [26–28]. Kakehi [16] pointed out that twin deformation in Ni-base single crystals loaded along [0 0 1] orientation occurs preferentially under compression since twinning shear on {1 1 1} planes in h1 1 2i direction is favored by compressive stress. The model proposed by Kolbe [24] assumes that the microtwins are formed by the passage of pairwise a/6 h1 1 2i Shockley partials on adjacent {1 1 1} planes. Experimental results of Kovarik et al. [25] support this assumption. In the ordered L12 structure of the c0 -phase having nominal stoichiometry Ni3Al, shearing of two a/6 h1 1 2i Shockley partials creates an energetically unfavorable pseudotwin with Al atoms as nearest neighbors. These pseudotwins should be converted into a true twin by a diffusion mediated reordering process. The reordering process is seen as rate-limiting for the microtwin deformation [25]. The temperature dependence of diffusion explains why microtwins form only at elevated temperatures. Thus, microtwins can form when compressive stress coincides with elevated temperatures, which is given in the case of OP loading. Since dwells allow for intensified diffusion reordering, their introduction will increase the extent of microtwinning as observed after OPTD tests. Thus, the present results are in agreement with the model of Kolbe [24] and the assumption that the diffusion reordering is rate limiting for the microtwin deformation. It has been reported that microtwinning results in easier creep deformation [16] and a reduced critical resolved shear stress [24]. From this can be inferred that the observed easier plastic deformation under compressive load at 750 °C resulting in a higher absolute mean stress in OPTD tests than in IPTD tests is produced by extensive microtwin formation. The cyclic increase of the relaxed stress along with the apparent change of relaxation mechanism during OPTD dwells (Fig. 4) may also be attributed to microtwin deformation. Localized plastic deformation along the {1 1 1} microtwin planes produces the surface deformation lines observed after OPTD loading whose orientation corresponds approximately to the trace of the {1 1 1} twin planes (Fig. 9c). Based on TMF OP experiments with dwell times on Ni-base single crystals, this was also concluded by Zhang et al. [29]. It is noticeable that during OPTD loading, the tensile cyclic yield stress decreased with increasing cycle number although under tensile loading at temperatures around 100 °C, microtwinning should not occur (Fig. 3). Microtwins along {1 1 1} planes apparently facilitate ordinary dislocation movement in {1 1 1}h1 1 0i slip system. Assuming that the loading axis is parallel to the initial [0 0 1] direction and the 1 1Þ, the twinned lattice would be loaded twin mirror plane is ð1 1 direction. The maximum Schmid factor for along the ½2 2 and ½2 2 1 is 0.408. {1 1 1}h1 1 0i slip for both loading along ½0 0 1 The Schmid factors of the cubic {1 0 0}h1 1 0i systems along ½0 0 1 1 are generally higher. According to is 0 and the values along ½2 2 the model of Lall et al. [15], this would result in an increased critical resolved shear stress in the twinned c0 lattice which is contrary to observations. Consequently, with this idealized consideration,
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the cyclic softening cannot be attributed to a change of crystal orientation due to the twins. Possibly, the necessary stress for dislocation slip across c–c0 interfaces is reduced when the interface is twinned. However, this needs to be investigated in further work. The cyclic hardening during IPC, OPC and particularly IPTD cycling (Figs. 2 and 3) can be explained by dislocation multiplication in the c-channels impeding further dislocation movement. The dislocation structure in the c-channels indicates that the same hardening effect takes also place in OPTD tests. However, in this case it is outweighed by softening due to microtwinning. In OPC tests, the amount of microtwins formed is apparently too low to cause macroscopic softening. 4.2. Damage mechanisms and lifetime behavior Although constant mechanical strain range has been applied for the four test types, the lifetimes differ significantly. The reason for this is that the dominant damage mechanisms vary with the TMF phase angle and may be enhanced by dwell times. Furthermore, additional damage mechanisms may operate during the dwells. In OP tests, the tensile half-cycle occurs at low temperature when few slip systems are active and the surface oxide layer has low ductility. Thus, cracks initiate easily at the surface oxides and propagate in low plasticity regime resulting in a flat and sharp morphology as shown in Fig. 9a. It has been frequently stated that deformation microtwins may provide a preferential crack path along which crack propagation accelerates [26–30]. Hence, principal cracks under approximately 45° to the loading axis and the transition of crack propagation from perpendicular to approximately 45° to the loading axis may be attributed to crack growth along the {1 1 1} microtwin planes. The introduction of dwell times leads to more frequent microtwinning which is accompanied by material softening. Because of the resulting higher plastic strain amplitude and more possible crack paths along microtwin planes, cracks grow more rapidly contributing to the observed decrease of fatigue lifetime due to dwells. Zhang et al. [30] found that the portion of crystallographic cracking along microtwins increases with the introduction of dwells in TMF-OP tests which supports this assumption. Since compressive stress acts during OPTD dwells, the cracks are not opened and oxidation of the crack tip area should have a minor effect on OPTD lifetime. However, surface oxidation and coarsening of superficial carbides during the dwells may promote the formation of preferential crack initiation sites. That way, the stage of crack initiation is shortened. Under IP loading, the tensile half-cycle occurs at elevated temperatures when the surface oxide layer has higher ductility. Therefore, the surface oxide layer has only minimum effect on crack initiation [22]. This is supported by the results of Gordon et al. [4] stating that for a given mechanical strain range, the critical oxide rupture thickness is significantly higher for IP than for OP loading. Similarly, coarse surface carbides developing in IPTD tests play a minor role for crack initiation. Consequently, IP specimens exhibit fewer surface cracks than OP specimens. Since at elevated temperatures more slip systems are active, IP cracks can propagate by repeated plastic blunting and re-sharpening which is the reason for their wavy morphology (Fig. 11a) [27]. The lifetime reduction due to dwells under IP loading may be attributed to various mechanisms. The increased internal damage after IPTD tests indicates that surface cracks linking up with internal cracks may accelerate their propagation rate. Creep crack propagation results of a Ni-base DS alloy at 760 °C suggest that cracks may grow substantially during the tensile dwells [10]. During IP dwells, the tensile loading opens up the cracks and the crack tips are exposed to ambient oxygen. Thus, oxidation induced formation of an embrittled zone [6,7] and c0 -dissolving [31] ahead of the crack tip during dwells may facilitate crack propagation in the following cycles.
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For both tests with and without dwells, IP lifetime was shorter than OP lifetime although OP loading results in a higher maximum stress. In case of OPTD tests, the plastic strain amplitude is also higher. The results are in agreement with reported shorter IP lives at high mechanical strain ranges for other Ni-base superalloys [2,32–34]. Considering that cracks initiate early in OP lifetime, it can be inferred that the average crack propagation rate under IP conditions is significantly higher than under comparable OP conditions. The observed numerous short surface cracks after OP loading compared to few long cracks after IP loading support this assumption. Assuming that cracks propagate predominantly in the tensile part of the cycle, this suggests that crack growth at high temperatures is much more rapid than at low temperatures. This may be related to the more homogeneous dislocation slip at high temperatures compared to planar slip at low temperatures [35]. He et al. [36] recently reported an increase of the growth rate of small cracks with temperature for a Ni-base DS alloy. They attributed this to material softening at elevated temperatures. At high temperatures, cracks are opened under IP loading and closed under OP loading. Hence, the role of oxidation induced embrittlement and c0 -dissolving ahead of the crack tip is presumably much more pronounced under IP loading which may be a further reason for higher IP crack propagation rates. Because of the complex interaction of fatigue, creep and oxidation damage combined with twin formation and stress redistributions, it is difficult to assess which damage mechanism governs the lifetime for the four test types, respectively. In order to quantify the contributions of the particular damage mechanisms, further work including tests with varying dwell time lengths and mechanical strain amplitudes is needed.
5. Conclusions The effects of 20 min dwell times on the deformation, damage and lifetime behavior of the directionally solidified Ni-base alloy Alloy 247 LC DS in IP and OP TMF tests with constant strain range of 1.5% and a temperature range of 100–750 °C has been studied. The following conclusions can be drawn: 1. Lifetime under IP loading is about one third of the lifetime under OP loading. For both phase angles, the introduction of 20 min dwells at 750 °C reduced the lifetime to about one sixth. 2. During OP cycling, microtwin deformation occurs which is intensified by dwell times. Microtwins apparently facilitate plastic deformation along {1 1 1} planes resulting in material softening. After IP loading, no microtwinning was observed. The crystallography/diffusion-based model for microtwinning introduced by Kolbe [24] could explain the dependence on phase angle and dwell times. 3. Fatigue crack propagation along microtwins in OP tests and formation of internal cracks at carbides and carbide/matrix interfaces in IP tests represent specific damage mechanisms identified in TMF of a directionally solidified Ni-base superalloy. Introducing dwells intensifies both mechanisms thus contributing to the observed lifetime reduction. Increased oxidation during dwells promoting superficial crack initiation under OP loading and embrittlement as well as c0 -dissolving ahead of the crack tip under IP loading may further reduce the lifetime. 4. During compressive dwell times of OP tests, the gauge length may exhibit transverse creep deformation under action of the force applied by the extensometer rods. That way the strain measurement is distorted which leads to an apparent increase of specimen stiffness. In order to avoid this, the use of self-supporting extensometers is suggested for any similar testing situation including compressive dwells at elevated temperatures.
Acknowledgements This work was supported by the Ministry of Education, Youth and Sports of the Czech Republic throughout the project No. CZ.1.07/2.3.00/30.0063 ‘‘Talented postdocs for scientific excellence in physics of materials‘‘. The financial support of the Karlsruhe House of Young Scientists (KHYS) for the stay abroad at the Institute of Physics of Materials (IPM) in Brno is gratefully acknowledged. Sincere thanks to Ludvik Kunz for the invitation to the research stay.
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