Cyclic deformation and lifetime of Alloy 617B during isothermal low cycle fatigue

International Journal of Fatigue 55 (2013) 126–135

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Cyclic deformation and lifetime of Alloy 617B during isothermal low cycle fatigue Gerhard Maier a,⇑, Hermann Riedel a, Christoph Somsen b a b

Fraunhofer Institute for Mechanics of Materials IWM, Woehlerstraße 11, 79108 Freiburg, Germany Ruhr-Universität Bochum RUB, Universitätsstraße 150, Geb. IA-1/130, 44801 Bochum, Germany

a r t i c l e

i n f o

Article history: Received 28 December 2012 Received in revised form 22 April 2013 Accepted 2 June 2013 Available online 10 June 2013 Keywords: Alloy 617 Precipitation Low cycle fatigue Cyclic deformation Fatigue life prediction

a b s t r a c t Isothermal low cycle fatigue tests are carried out on the nickel-base Alloy 617B in the solution-annealed, stabilized and long-term aged conditions at temperatures between room temperature and 900 °C. In addition, fatigue microcrack growth is measured using the replica technique. Transmission electron microscopy studies suggest that the observed differences in cyclic hardening between the different heat treatments result from the precipitation of fine carbides. Scanning electron microscope observations indicate a change in fracture mode for the solution-annealed and long-term aged material with temperature. The Chaboche model is able to describe the time and temperature dependent cyclic plasticity of the three material conditions. The measured lifetimes and crack growth rates can be described using a fracture mechanics based lifetime model. However, the data for room temperature and for temperatures above 400 °C fall into two different scatter bands due to differences in crack growth rates. Ó 2013 Elsevier Ltd. All rights reserved.

1. Introduction Advanced ultra super critical (A-USC) power plants with steam inlet temperatures of approx. 700 °C are the most promising technology to improve thermal efficiency and therefore reduce CO2 emissions of fossil fuel power plants. Additionally, new A-USC power plants need to be more flexible to compensate the fluctuating output of renewable energy sources. Thus, power plant materials require not only high creep strength, but also high resistance to low cycle fatigue (LCF). The nickel-base Alloy 617B, which is a modified version of Alloy 617 [1,2], has been selected as a candidate material for key components of the A-USC power plants. Several research programs are concerned with evaluating the long-term properties of Alloy 617B. Hence, there is a need for LCF data along with adjusted material models that are able to predict the fatigue behavior under service conditions. In previous works, Maier et al. [3,4] already published experimental results of Alloy 617B in different heat treatments together with suitable models for the description of cyclic deformation and lifetime: In [3], a short overview of the time and temperature dependent cyclic plasticity and fatigue crack growth of Alloy 617B is presented. In [4], the cyclic deformation and lifetime of Alloy 617B under non-isothermal loading are discussed in detail. Furthermore, results of transmission electron microscopy ⇑ Corresponding author. Tel.: +49 761 5142 431; fax: +49 761 5142 510. E-mail addresses: [email protected] (G. Maier), [email protected] (H. Riedel), [email protected] (C. Somsen). 0142-1123/$ - see front matter Ó 2013 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.ijfatigue.2013.06.001

investigations, crack growth measurements at non-isothermal loading and fractographic investigations of tested specimens are presented. The aim of this work is to give a detailed description of the isothermal LCF properties of Alloy 617B in the solution-annealed, stabilized and long-term aged conditions. The work includes crack growth measurements and fractographic investigations. In addition, the adjustment of the already proposed models for cyclic viscoplasticity and fatigue crack growth to Alloy 617B is presented in more detail in this work. 2. Material and experimental The solid-solution and carbide-hardened material Alloy 617B (NiCr23Co12Mo B, NicroferÒ5520 CoB) is a special version of the conventional Alloy 617 with improved creep properties in the temperature range from 650 °C to 720 °C. The chemical composition of the investigated material is given in Table 2.1. The material was supplied by ThyssenKrupp VDM (now Outokumpu VDM) as 30 mm thick hot rolled sheets in three different conditions: Solution-annealed (LG, 1175 °C/1 h/water), stabilized (SGH, 1175 °C/1 h/water + 980 °C/3 h/air) and long-term aged (SGA, 1175 °C/1 h/water + 980 °C/3 h/air + 700 °C/1 yr/air). The long-term aged condition represents an intermediate service-like condition with an aged microstructure. The stabilizing heat treatment is applied in order to reduce the susceptibility to intergranular cracking after short-term service [5], whereas the alloy is conventionally used in the solution-annealed condition [6]. The

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Table 2.1 Chemical composition of Alloy 617B as-received (in wt%). Ni Bal.

Cr 22.2

Co 11.6

Mo 8.7

Al 1.0

C 0.06

Fe 1.0

Ti 0.35

Si 0.14

Mn 0.06

B 0.004

room temperature tensile properties of the different material conditions are given in Table 2.2. The tensile tests were conducted with a nominal strain rate of 104 1/s. Fig. 2.1a shows a light optical micrograph of the LG material with a mean grain diameter of 135 lm and several twin boundaries. Carbides (grey) and primary titanium(carbo-)nitrides (orange) were found in all heat treatments. Fig. 2.1b shows the light optical micrograph of the SGH material with secondary carbides at grain boundaries and inside the grain. In order to investigate the microstructure of the three material conditions more precisely, transmission electron microscopy (TEM) was used. The investigations, presented in detail in [4], reveal that the precipitation morphology of secondary carbides depends on the preceding heat treatment: For the LG material M23C6 precipitates can be detected only at large angle grain boundaries, whereas for the SGH and SGA material, M23C6 precipitates can be detected inside the grains, at large angle grain boundaries and occasionally also at twin boundaries. The size of the carbides varies between only a few tenths of nanometers in the LG material and up to 500 nm in the SGH and SGA material. Furthermore, for the SGA material also c0 precipitates can be detected inside the matrix with a size of around 50–100 nm. MC and M6C carbides could not be found in any of the samples. Fatigue tests were performed on round specimens with a circular cross section with 7 mm gauge diameter using an electromechanical testing machine from Instron Ltd. Elongation was measured with a high temperature Maytec capacitive extensometer with 10 mm gauge length. The specimens were heated inductively. The temperature was measured and controlled with three Ni–CrNi thermocouples positioned in the center, the upper and lower part of the gauge length to monitor the homogeneity of the temperature distribution. At the beginning of an experiment, each specimen was heated up to test temperature at zero stress. Complex LCF (CLCF) tests were performed under fully reversed strain control (with strain ratio of 1) at room temperature (RT), 400 °C, 600 °C, 700 °C, 800 °C and 900 °C with total strain amplitudes varying from 0.2% to 0.9%. Each CLCF experiment consisted of an aperiodic and a periodic part. In the aperiodic part, different strain rates and hold times up to 30 min are used. In the periodic part, the specimen was cyclically loaded with a nominal strain rate of 103 1/s and constant strain amplitude until failure [7]. Thus, effects of the loading rate, relaxation and cyclic material behavior can be determined in a single experiment. Fig. 2.2 shows schematically the loading history applied in CLCF tests. In all experiments, the number of cycles to failure Nf was determined by a 5% drop of the maximum tensile load. Furthermore, the influence of a thermal and thermo-mechanical loading on the evolution of the microstructure of the LG and SGH material at 700 °C is studied using TEM. To this end, specimens were either aged or thermo-mechanically loaded with total strain amplitude of 0.4% and nominal strain rate of 103 1/s for 3 h, respectively. TEM was performed using a FEI Tecnai F20G2 operating at 200 kV, equipped with a scanning unit (STEM) and high-an-

gle annular dark-field (HAADF)-detector. Thin foils were prepared from 3 mm disks, which were ground down to a final thickness of about 120 lm. Electron transparent regions were obtained by electro-polishing using a standard double-jet procedure (TenuPol-5, Struers) with an electrolyte consisting of 5% perchloric acid and 95% acetic acid at 15 °C and a voltage of 55 V. Crack growth measurements of surface cracks were performed at RT and 700 °C with SGH material using the replica technique [4]. Fracture surfaces were examined by scanning electron microscopy (SEM) using a CAMSCAN CS 24 to determine the crack initiation and propagation modes. 3. Time and temperature dependent cyclic plasticity 3.1. Experimental results Figs. 3.1–3.3 show the stress response to the aperiodic and periodic parts of the CLCF tests for all temperatures and selected strain amplitudes. Stress relaxation during hold times is first visible at 700 °C. A pronounced strain rate dependency is observed only at 800 °C and 900 °C. At RT, all specimens exhibit a relatively short period of cyclic hardening in the periodic loading stage, followed by a period of nearly stable peak stress with minor cyclic softening. Toward the end of the tests, the load decreases rapidly indicating the formation of macrocracks. In general, the magnitude of peak stresses depends on the preceding heat treatment: At RT, LG material shows the lowest peak stresses, SGA material the highest. At higher temperatures the alloy undergoes continuous cyclic hardening with a slightly higher degree of hardening for the LG material than for the SGH and SGA materials. This effect increases with temperatures up to 800 °C. At 800 °C, the LG material shows higher stable peak stresses in the periodic part of the experiment compared to SGH and SGA, and cyclic hardening for the LG and SGH material seems to saturate already after few loading cycles. In contrast, the stress response of the SGA material exhibits softening already during the aperiodic part of the experiment. However, the decreasing peak stresses during the periodic part of the experiment most probably result from a continuous necking of the fatigue specimen as a consequence of difficulties in strain control due to severe dynamic strain aging (Portevin-Le Chatelier effect). At 900 °C, there is no difference in material response of the LG, SGH and SGA material, whereas no tests were performed with SGA material. Furthermore, for temperatures from 400 °C to 900 °C, dynamic strain aging is observed by serrated stress curves. For better illustration of the stress response the abrupt stress drops due to dynamic strain aging have been deleted in Figs. 3.1–3.3. 3.2. Chaboche model The viscoplastic Chaboche model [8,9] is used to describe the time and rate dependent deformation at high temperatures. The total strain rate of the isothermal Chaboche model (for uniaxial stress states) is given by

e_ tot ¼ e_ v p þ

r_

ð1Þ

E

with the viscoplastic strain rate e_ v p , the stress rate r_ and the temperature dependent Young’s modulus E. The viscoplastic flow rule and the equivalent viscoplastic strain rate are given by

Table 2.2 Room temperature tensile properties of LG, SGH and SGA material.

LG SGH SGA

Ultimate strength (MPa)

Yield strength (MPa)

Total elongation (%)

Reduction in area (%)

753 804 902

313 316 439

69.4 56.1 36.0

56.1 43.3 33.0

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Fig. 2.1. Micrographs of (a) LG material, etched in Adler reagent: mean grain diameter 135 lm, and (b) SGH material, etched in Beraha 2 reagent: twinned grains, titanium(carbo-)nitrides (orange) and secondary carbides (white) at grain boundaries and inside the grain. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)

e_ v p ¼ p_

@U where p_ ¼ @r

 n

U K

:

ð2Þ

K and n are temperature dependent model parameters. U is the yield function with kinematic and isotropic hardening:

U ¼ jr  Xj  ðy0 þ RÞ:

ð3Þ

X is the backstress, y0 is the initial yield stress and R the isotropic hardening function

R ¼ Q 1 ð1  exp½bpÞ

ð4Þ

with the temperature dependent material parameters Q1 and b. The backstress is additively decomposed into two parts, X = X1 + X2, each following the evolution equation [8,9]:

_ i  Ri X i X_ i ¼ C i e_ v p  ci ui pX

ð5Þ

Fig. 2.2. Prescribed strain in CLCF tests with logarithmic time scale.

Fig. 3.1. Comparison of experimental data and model prediction of CLCF tests with LG, SGH and SGA material at RT (upper figures) and 400 °C (lower figures) with ea,t = 0.5%, left: total stress response, where in the periodic part only peak stresses are shown, right: hysteresis loops at half lifetime.

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Fig. 3.2. Comparison of experimental data and model prediction of CLCF tests with LG, SGH and SGA material at 600 °C (upper figures) and 700 °C (lower figures) with ea,t = 0.4%, left: total stress response, where in the periodic part only peak stresses are shown, right: hysteresis loops at half lifetime.

Fig. 3.3. Comparison of experimental data and model prediction of CLCF tests with LG, SGH and SGA material at 800 °C (upper figures) and 900 °C (lower figures) with ea,t = 0.3%, left: total stress response, where in the periodic part only peak stresses are shown, right: hysteresis loops at half lifetime.

The first two terms of Eq. (5) are the Armstrong and Frederick [10] kinematic hardening law. The third term represents the static recovery term, which was proposed by Chaboche [9]. In Eq. (5), ui is used to describe cyclic hardening and softening

Ci, ci, Ri , u1,i and xi are the temperature dependent model parameters.

ui ¼ u1;i þ ð1  u1;i Þ exp½xi p:

The calibration of the Chaboche model parameters was carried out at first on experimental data from isothermal tests with SGH

ð6Þ

3.3. Adjustment of the model parameters

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material. The adjustment was accomplished by minimizing the error square between experimental and calculated data for each test temperature separately. To account for the differences in viscoplastic material behavior of the LG and SGA material in comparison to the SGH material, the isotropic hardening parameters Q1, b, y0 were identified for each material individually, whereas all other material parameters were kept constant for all three heat treatments. Figs. 3.1–3.3 show the comparison of experimental data with the model response for the different materials and selected loading conditions. 4. Fatigue crack growth

Fig. 4.1 shows the correlation between the damage parameter dnZD/rcy and Nf for all conducted CLCF tests. The dashed lines in Fig. 4.1 represent the prediction according to Eq. (10). Apparently the tests with temperatures above 400 °C fall into a common scatter band around the theoretical line with A = 6. Room temperature data lie near a line with A = 48, and the 400 °C data lie in between. The stabilizing heat treatment extends the lifetime only at 800 °C and 900 °C whereas at lower temperatures LG specimens exhibit longer fatigue lives than SGH specimens. In general, fatigue life of the long-term aged material is reduced for temperatures lower than 600 °C compared to LG and SGH material. However, at service temperature of 700 °C, fatigue life is not affected by the preceding heat treatment.

4.1. Mechanism-based lifetime model 4.2. Crack growth measurements The mechanism-based approach to model the time and temperature dependent microcrack growth in LCF tests assumes that the crack advance per loading cycle, da/dN, is proportional to the crack-tip opening displacement DCTOD:

da ¼ bDCTOD; dN

ð7Þ

where b is a proportionality constant. DCTOD is analytically estimated using the cyclic J-integral Z resulting in DCTOD = dnZ/rcy, where dn is given by the polynomial dn ¼ 0:78627  3:41692n0 0 þ6:11945n02  4:2227n03 derived from [11] with n being the Ramberg–Osgood hardening exponent. rcy is the cyclic yield stress defined as the 0.2%-offset stress with respect to the point of load reversal. The cyclic J-integral can be approximated by the sum of an elastic and plastic contribution

Z¼

1:45

! Dr2eff 2:4 þ pffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi DrDev p a ¼ Z D a: E 1 þ 3n0

ð8Þ

The damage parameter ZD was proposed by Heitmann et al. [12] and was derived for semicircular surface cracks and rate-independent materials. a is the crack length and E is the Young’s modulus, Dr and Devp the stress and plastic strain range. Crack closure is taken into account by the use of the effective stress range

Dreff ¼

ð1  rop =rmax Þ Dr 1R

ð9Þ

in the elastic part of ZD. rop is the crack opening stress determined by the empirical crack opening stress equation by Newman [13]. rmax is the maximum stress and R the ratio of minimum and maximum stress. By integration of (7), the number of cycles to failure Nf can be calculated by

 1 lnðaf =a0 Þ ZD Nf ¼ A dn with A ¼ : rcy b

ð10Þ

The factor A includes the crack length at failure af and the initial crack length a0. For the calculation of the damage parameter, the stress and 0 strain values, E, and n are determined from stabilized hysteresis loops at half lifetime. To this end, the Ramberg–Osgood relation

de ¼

0 0 dr þ 0:002ðrcy Þ1=n ðdrÞ1=n E

ð11Þ

is adjusted to the rising branch of the hysteresis. de and dr are the strain and stress difference to the lower point of load reversal. The results of the hysteresis analysis of the CLCF tests from Figs. 3.1–3.3 are summarized in Table 4.1. To investigate the scatter of hysteresis parameters and fatigue life, CLCF tests with SGH material have been performed under identical conditions. The respective results of the hysteresis analysis at half lifetime are summarized in Table 4.2.

Fatigue crack growth of individual cracks was measured at RT and 700 °C, using the replica technique. Both tests were carried out on SGH material with total strain amplitude of 0.5%. Fig. 4.2 shows the measured half surface crack lengths as a function of the number of cycles. Surface crack growth is characterized by periods of slow growth interrupted by jumps, which correspond to coalescence events of surface microcracks. The jumps are more pronounced at RT than at 700 °C. After coalescence, crack growth at the surface is delayed. An identical crack growth behavior could also be found at non-isothermal loading between 50 °C and 700 °C [4]. In Fig. 4.2, the theoretical growth law a = a0 exp(bdnZD/rcyN) derived from Eq. (7) is fitted to the measured curves with the adjustable parameter b = 0.09805 at RT and b = 0.80745 at 700 °C. Hence the difference in lifetime by a factor of eight is reflected by the same difference in crack growth rate. 5. Microstructural investigations 5.1. TEM observations Figs. 5.1–5.3 show results of TEM investigations of LG and SGH specimens aged for 3 h at 700 °C both with and without LCF load (ea,t = 0.4%, e_ ¼ 103 1=s). Thus, the influence of a purely thermal and superimposed mechanical loading on the evolution of the microstructure can be studied. The microstructural observations can also be used to explain the measured differences in the cyclic stress response of the LG and SGH material during LCF loading at 700 °C, see Fig. 3.2. Fig. 5.1 shows typical TEM images of the LG specimen after the 3 h aging treatment at 700 °C. No relevant changes can be found in the microstructure compared to the solution-annealed reference condition: Chromium-rich carbides of the type M23C6 can be found only at large-angle grain boundaries, whereas no MC, M6C or c0 precipitates are detectable. In Fig. 5.2 typical TEM images of the SGH specimen after 3 h aging at 700 °C are presented. As in the stabilized reference condition, coarse chromium-rich M23C6 carbides can be detected at grain boundaries and inside the matrix, varying between 100 nm and 500 nm in size. As for the LG material, no evidence of new (fine) carbide precipitations could be found inside the matrix. Fig. 5.3 shows results of LG and SGH specimens LCF loaded at 700 °C with total strain amplitude of 0.4% and nominal strain rate of 103 1/s for 3 h. In Fig. 5.3a a STEM–HAADF image of a typical region inside a grain of the LG material is shown, where fine M23C6 precipitates can be detected on dislocations. Due to the fact that their size is less than 20 nm they could not be captured by EDX measurements but the presence is proven by their characteristic reflections in the respective diffraction pattern presented in

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G. Maier et al. / International Journal of Fatigue 55 (2013) 126–135 Table 4.1 Results of the hysteresis analysis at Nf/2 of selected CLCF tests.

LG SGH SGA LG SGH SGA LG SGH SGA LG SGH SGA LG SGH SGA LG SGH SGA

Temp. (°C)

De (–)

Nf (–)

rmax/rmin (MPa)

E (GPa)

rcy (MPa)

n0 (–)

ZD (MPa)

RT RT RT 400 400 400 600 600 600 700 700 700 800 800 800 900 900 900

0.01 0.01 0.01 0.01 0.01 0.01 0.008 0.008 0.008 0.008 0.008 0.008 0.006 0.006 0.006 0.006 0.006 0.006

8330 6235 6227 3260 2753 2308 2315 1903 1741 1072 965 1525 1128 1463 1344 946 1216 913

463/475 476/488 529/534 490/506 473/492 525/518 507/534 471/486 485/495 504/523 416/430 481/489 320/327 281/281 297/300 228/235 226/228 204/195

222 220 210 198 198 195 188 188 185 183 181 162 178 172 172 163 160 156

797 810 884 879 836 924 1020 919 945 1010 794 955 635 544 582 442 415 378

0.1506 0.1685 0.1773 0.1346 0.1358 0.1350 0.1029 0.1122 0.1181 0.0925 0.0909 0.0995 0.0706 0.0716 0.0752 0.0651 0.0627 0.0761

14.130 14.215 14.997 13.942 13.771 14.418 9.020 9.043 8.985 11.250 8.759 8.733 4.881 4.622 4.674 4.108 4.165 3.756

Table 4.2 Results of hysteresis analysis at Nf/2 of CLCF tests under identical loading conditions.

SGH SGH SGH SGH

0

Temp. (°C)

De (–)

Nf (–)

rmax/rmin (MPa)

E (GPa)

rcy (MPa)

n (–)

ZD (MPa)

700 700 700 700

0.006 0.006 0.006 0.006

2172 1611 1906 2116

396/408 387/403 400/410 393/408

178 182 182 183

825 800 825 812

0.0872 0.0821 0.0833 0.0886

4.722 4.895 4.889 4.890

Fig. 4.1. Damage parameter dnZD/rcy over number of cycles to failure Nf for all tests with LG, SGH and SGA material. The dashed lines represent the prediction according to Eq. (10). RT data lie near a line with A = 48, experiments with test temperature >400 °C lie near a line with A = 6.

Fig. 5.3b. Fig. 5.3c shows a STEM–HAADF image of a typical region inside a grain of the SGH material where no coarse carbides are present. In the respective diffraction pattern shown in Fig. 5.3d no evidence of fine carbide precipitates could be found inside the SGH material.

5.2. Fractography After fatigue testing, fracture surfaces from all test temperatures were investigated using SEM. The analysis of the fracture surfaces of RT specimens indicate an identical fracture mode for the LG and SGH material with purely transgranular crack initiation and propagation with transgranular facets and fatigue striations. Fig. 5.4a shows the region of nucleation and early propagation of

Fig. 4.2. Evolution of the crack length at RT and 700 °C. Lines with symbols: measured at the specimen surface; solid lines: calculated from Eq. (7) with adjusted b. Surface crack growth is characterized by coalescence of individual microcracks.

fatigue cracks typical for the LG and SGH material at RT. Crack initiation in the SGA material also was transgranular and occurred in a flat, faceted manner, identical to the LG and SGH material. For longer cracks, isolated grain boundaries (GB) appear within the transgranular fracture surface, see Fig. 5.4b. With increasing temperature, the number of isolated broken grain boundaries on the fracture surface of the SGA material reduces significantly. At 700 °C fracture mode is identical for all three heat treatments: Crack nucleation and propagation are purely transgranular, Fig. 5.5a, residual fracture is intergranular, Fig. 5.5b. At 800 °C, fatigue crack nucleation in the LG material occasionally occurs at grain boundaries, whereas crack propagation is still purely transgranular. In contrast, SGH and SGA material still exhibit purely transgranular crack nucleation and propagation. At 900 °C, fatigue crack nucleation and very early crack propagation

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Fig. 5.1. TEM images of LG specimen after aging at 700 °C/3 h, (a) STEM–HAADF micrograph of a large angle grain boundary, (b) TEM diffraction pattern in [1 0 0] zone direction of a typical region inside a grain: no additional reflections beside that of the fcc matrix are present, (c) EDX of position 1 in (a) detecting an increased chromium content indicating carbide precipitations, and (d) EDX of position 2 in (a) detecting matrix material.

in the LG material occur along grain boundaries, although after two to three grains, crack propagation reverts to a transgranular mode, Fig. 5.6a. In contrast, in the SGH material crack nucleation and propagation still occur purely transgranular, Fig. 5.6b. For longer crack lengths, fatigue striations were observed on the fracture surfaces, interspersed with isolated transgranular facets. At both temperatures, residual fracture occurs in an intergranular fracture mode. With SGA material no SEM investigations were performed at 900 °C. 6. Discussion 6.1. Microstructural investigations From TEM observations of specimens in the as-received conditions, carbides of the type M23C6 could be detected in all of the three heat treatments. However, size and number of carbides depend on the preceding heat treatment. In the LG material only small carbides can be detected at large angle grain boundaries, whereas after the stabilizing heat treatment at 980 °C/3 h, several large carbides could also be found at grain boundaries and near dislocations inside the grain. Carbide sizes vary between a few nanometers in the LG material and up to 500 nm after the stabilizing heat treatment. In contrast to the carbides, precipitates of the c0 phase were only found after long-term aging at 700 °C (SGA material). This is in agreement with the work of Wu [14] who found c0 precipitates in specimens aged at 704 °C for 43,100 h and 65,600 h. He also stated that c0 phase could be observed after 80 h aging time at 700 °C in previous works. In the work of Gariboldi et al. [15], numerous c0 precipitates could also be detected after aging and creep exposure at 700 °C for 34,000 h.

The measured differences in initial yield strength of the SGA material in comparison to the LG and SGH material (Table 2.2) can therefore directly be related to the precipitation of the c0 phase during the 700 °C/1 yr aging. On the other side, due to the resulting carbide distribution and size, the stabilizing heat treatment does not affect the yield strength of the material considerably. Furthermore, the evolution of the microstructure of the LG material at 700 °C strongly depends on the applied loading: After thermal aging of the LG material for 3 h at 700 °C no differences in the microstructure can be observed in comparison to the solution-annealed reference condition. However, with superimposed fatigue loading numerous fine M23C6 carbides were precipitated after 3 h inside the grains. These fine precipitates with a size <20 nm lead to a substantial cyclic hardening of the LG material, Fig. 5.3. After stabilizing heat treatment, no fine carbides can be precipitated in the matrix during cyclic loading, likely because most of the carbon is already bound in large carbides. It is assumed, that the precipitation of fine carbides is due to the fact that the number of dislocations is drastically increased by the superimposed mechanical loading, whereat the dislocations serve as precipitation sites. Additionally, precipitation kinetics is accelerated due to an enhanced diffusion of carbide forming elements along dislocations (pipe diffusion). Enhanced nucleation and growth of precipitates in solution-annealed Alloy 617 under fatigue loading compared with isothermal aging for equivalent time durations and temperatures was already observed by Bhanu Sankara Rao et al. [16]. From Figs. 5.1a and 5.2a it becomes obvious that the precipitation of carbides might also generate dislocations. Mankins et al. [17] found that the growth of existing carbides would create dislocations on which smaller carbides could precipitate during

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Fig. 5.2. TEM images of SGH specimen after aging at 700 °C/3 h, (a) STEM–HAADF micrograph, (b) TEM diffraction pattern in [1 0 0] zone direction of a typical region inside a grain, where no large carbides are present: no additional reflections beside that of the fcc matrix can be detected, (c) and (d) EDX of positions 1 and 2 in (a) both detecting an increased chromium content indicating carbide precipitations.

subsequent temperature exposure. As a result, after solutionannealing, nondissolved precipitates can already grow and create new nucleation sites, as the alloy is cooled down from annealing temperature. 6.2. Cyclic plasticity and lifetime – experiments and models Viscoplastic deformations measured in the isothermal CLCF tests from Figs. 5.1–5.3 depend on the heat treatment and thus on the microstructure. As discussed above, the pronounced cyclic hardening of the solution-annealed material at 700 °C is due to the precipitation of fine M23C6 carbides. This was also observed by Mankins et al. [17] in solution-annealed Alloy 617 at 760 °C. Further, Burke and Beck [18] identified a much denser precipitation-stabilized dislocation structure as the reason for markedly higher cyclic hardening at 760 °C compared to higher temperatures. Bhanu Sankara Rao et al. [16] attributed the substantial cyclic hardening at 750 °C to the cumulative effects of dislocation generation with their mutual interaction and to the immobilization of dislocations by carbide precipitations. The Chaboche model describes several aspects of the viscoplastic material behavior well. Interestingly, the differences in viscoplastic material behavior between the three heat treatments (solution-annealed, stabilized and long-term aged) can be captured by adjusting the isotropic hardening parameters only, rather than the whole parameter set. The measured fatigue lives are almost identical for all three heat treatments: The stabilizing heat treatment extends the lifetime in isothermal LCF tests only at 800 °C and 900 °C, whereas at 700 °C, lifetime is almost independent of the preceding heat treatment.

The stabilizing heat treatment improves the creep behavior of the material and is therefore only effective at temperatures where creep deformation occurs. The effect of a stabilizing heat treatment on the creep life was already shown at 950 °C in [19], where the creep ductility could be considerably increased by a preceding heat treatment at 900 °C/1000 h. Investigations on the lifetime scatter show insignificant differences in measured lifetimes in case of LCF loading at 700 °C. The applied lifetime model based on cyclic crack-tip opening and the cyclic J-integral for microcrack growth allows a unified description of the lifetimes in all LCF tests above 400 °C, while the room temperature data fall into a separate scatter band with about eight times longer lifetimes. Measurements of the crack length evolution by the replica technique show that microcracks at 700 °C grow by about a factor eight faster than at room temperature. This is consistent with the measured difference in lifetimes. The lifetime model, on the other hand, would predict a common scatter band for all temperatures. From crack length measurements, it also becomes apparent, that microcracks growth is characterized by the coalescence of surface cracks. It is assumed, that the surface crack growth is delayed until the crack has established its approximately semicircular equilibrium shape by preferential growth in depth direction. No reason for the difference in lifetime at RT and high temperatures has been found by the fractographic investigations. For all temperatures and material conditions, fatigue crack initiation was transgranular and occurred in a flat, faceted manner, except for the solution-annealed material where at 800 °C and 900 °C crack initiation was intergranular. At 900 °C, not only the crack initiation but also the very early crack propagation is intergranular.

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Fig. 5.3. Stress response and TEM images of LG and SGH specimens cycled at 700 °C with ea,t = 0.4 % and e_ = 103 1/s for 3 h, (a) STEM–HAADF micrograph of LG material with dislocations that appear with white contrast, (b) TEM diffraction pattern in [1 0 0] zone direction of (a) with reflections of the fcc matrix and M23C6 carbides, (c) STEM–HAADF micrograph of SGH material of a region where no coarse carbides are present, and (d) TEM diffraction pattern in [1 0 0] zone direction of (c) with reflections of the fcc matrix only.

Fig. 5.4. SEM fractographs of fatigue crack initiation and early propagation at RT (a) image of SGH material with purely transgranular crack initiation and propagation typical for LG and SGH material, and (b) image of SGA material with isolated grain boundaries (GB) within the generally transgranular fracture surface.

Fig. 5.5. SEM fractographs of SGA material at 700 °C with typical features of fracture surfaces for LG, SGH and SGA material (a) fatigue crack initiation and early propagation occur in purely transgranular mode and (b) transition zone to residual fracture region with transgranular crack propagation (upper area) and intergranular residual fracture (lower area).

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Fig. 5.6. SEM images from fracture surfaces tested at 900 °C (a) typical crack initiation zone in LG material with intergranular initiation (left) and transgranular propagation (right) and (b) typical crack initiation zone in SGH material with purely transgranular crack initiation and propagation and intergranular residual fracture (R).

This is in agreement with the observations by Burke and Beck [18], who performed LCF tests with Alloy 617 in the solution-annealed condition at 871 °C. On the other hand, the stabilizing heat treatment prevents intergranular fatigue crack initiation and propagation. The observed transgranular facets are typical for nickel and cobalt alloys and occur as fatigue cracks tend to follow glide planes causing the impression of cleaved grains. For the tests performed in this work crack propagation occurs in a transgranular mode, except for the low temperature tests with the SGA material which show isolated intergranular facets on the otherwise transgranular fracture surface. 7. Conclusion Isothermal LCF properties of Alloy 617B in three different heat treatments have been investigated. The high initial yield strength of the SGA material results from the precipitation of the c0 phase during long-term aging at 700 °C. The pronounced cyclic hardening at 700 °C in the solution-annealed material condition is due to the precipitation of fine M23C6 carbides near dislocations. Nucleation and growth of these carbides is accelerated under cyclic loading and could not be observed in the stabilized and long-term aged material. Dislocation induced c0 precipitation could not be observed either in the LG material or in the SGH material. Thus, the formation of the c0 phase depends rather on the exposure time than on the applied loading. Suitable material models are proposed for the description of the viscoplastic deformation and lifetime. At the present time, the evolution of the microstructure is not taken into account by the deformation model. Furthermore, the difference in lifetime between room temperature and high temperatures under identical loading conditions can be ascribed to different crack growth rates. The performed experiments indicate that low cycle fatigue life is almost independent of the preceding heat treatment. However, the stabilizing heat treatment reduces the susceptibility to intergranular crack initiation under cyclic loading at high temperatures. Acknowledgments The authors want to acknowledge their thanks to Prof. G. Eggeler for supporting the TEM investigations and to Dr. J. Klöwer and Dr. R. Mohrmann for fruitful discussions. Furthermore, financial

support of ThyssenKrupp VDM GmbH and RWE Power AG is gratefully acknowledged. References [1] Final report, project No. A130, AVIF, contract No. 0327062, BMWA (MARCKODE2); 2005. [2] Blum R, Vanstone RW. Materials development for boilers and steam turbines operating at 700 C, in materials for advanced power engineering 2006. In: Proc. of the 8th liege conference; 2006. p. 41–60. [3] Maier G, Riedel H, Seifert T, Klöwer J, Mohrmann R. Time and temperature dependent cyclic plasticity and fatigue crack growth of the nickel-base Alloy 617B – experiments and models. Adv Mater Res 2011; 278:369–74. [4] Maier G, Riedel H, Nieweg B, Somsen C, Eggeler G, Klöwer J, et al. Cyclic deformation and lifetime of Alloy 617B during thermomechanical fatigue. Mater High Temp 2013;30(1). [5] van Wortel H. Control of relaxation cracking in austenitic high temperature components. In: Proceedings of the CORROSION 2007 conference, 2007, Paper No. 07423; 2007. [6] ThyssenKrupp VDM-Werkstoffdatenblatt Nr. 4119, Nicrofer 5520Co – alloy 617; 2005. [7] T. Seifert. In: Borsutzki M, Geisler S, editors. Werkstoffprüfung 2006 – Fortschritte der Kennwertermittlung für Forschung und, Praxis; 2006. p. 409. [8] Chaboche JL. Time-independent constitutive theories for cyclic plasticity. Int J Plast 1986;2:149–88. [9] Chaboche JL. Constitutive equations for cyclic plasticity and cyclic viscoplasticity. Int J Plast 1989;5:247–302. [10] Armstrong PJ, Frederick CO. A mathematical representation of the multiaxial bauschinger effect. Technical Report RD/B/N 731, C.E.G.B.; 1966. [11] Shih CF. Relationships between the J-intergral and the crack opening displacement for stationary and extending cracks. J Mech Phys Solids 1981;29(4):305–26. [12] Heitmann HH, Vehoff H, Neumann P. Life prediction for random load fatigue based on the growth behavior of microcracks. In: Valluri SR et al., editors. Advances in fracture research 1984 – proceedings of ICF6, vol. 5. New York: Pergamon Press, Oxford; 1984. p. 3599–606. [13] Newman JC. A crack opening stress equation for fatigue crack growth. Int J Fract 1984;24:131–5. [14] Wu Q. Microstructural evolution in advanced boiler materials for ultrasupercritical coal power plants. Dissertation 2006. [15] Gariboldi E, Cabibbo M, Spigarelli S, Ripamonti D. Investigation on precipitation phenomena of Ni–22Cr–12Co–9Mo alloy aged and crept at high temperature. Int J Pres Ves Piping 2008;85:63–71. [16] Bhanu Sankara Rao K, Schiffers H, Schuster H, Nickel H. Influence of time and temperature dependent processes on strain controlled low cycle fatigue behavior of Alloy 617. Metall Trans 1988;19A:359–71. [17] Mankins WL, Hosier JC, Bassford TH. Microstructure and phase stability of inconel Alloy 617. Metall Trans 1974;5:2579–90. [18] Burke MA, Beck CG. The high temperature low cycle fatigue behavior of the nickel base alloy IN-617. Metall Trans 1984;15A:661–70. [19] von der Gracht R, Ennis PJ, Czyrska-Filemonowicz A, Quadakkers WJ, Schuster H, Nickel H. Einfluß der Carbidausscheidungen auf das Kriechverhalten von Hochtemperaturlegierungen. Bericht der Kernforschungsanlage Jülich  Nr. 2211; 1988. p. 60.