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Low-cyclic fatigue behavior of modified 9Cr–1Mo steel at elevated temperature
Materials Science & Engineering A 604 (2014) 196–206
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Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Low-cyclic fatigue behavior of modified 9Cr–1Mo steel at elevated temperature Krishna Guguloth a,n, S. Sivaprasad a, D. Chakrabarti b, S. Tarafder a a b
CSIR-National Metallurgical laboratory, Material Science and Technology Division, Jamshedpur 831007, India Department of Metallurgical and Materials Engineering, Indian Institute of Technology, Kharagpur 721302, India
art ic l e i nf o
a b s t r a c t
Article history: Received 7 January 2014 Received in revised form 18 February 2014 Accepted 21 February 2014 Available online 2 March 2014
The low-cycle fatigue behavior of indigenously developed modified 9Cr–1Mo steel has been evaluated using a constant strain rate (1 10 3 s 1) at ambient temperature (25 1C) and at elevated temperatures (500–600 1C) over the strain amplitudes varying between 70.7% and 71.2%. Cyclic stress response showed a gradual softening regime that ended in a stress plateau until complete failure of the specimens. The estimated fatigue life decreased with the increase in test temperature. The effect of temperature on fatigue life was more pronounced at lower strain amplitudes. The cyclic deformation behavior at different temperatures has been analyzed from hysteresis loop and also in view of the changes taking place in dislocation structure and dislocation–precipitation interaction. Evaluation of low-cycle fatigue properties of modified 9Cr–1Mo steel over a range of test temperature can help in designing components for in-core applications in fast breeder reactors and in super heaters for nuclear power plants. & 2014 Elsevier B.V. All rights reserved.
Keywords: Low-cycle fatigue Modified 9Cr–1Mo steel Hysteresis loop Masing behavior Dislocation Precipitation
1. Introduction Martensitic steels with 9–12 wt% Cr has been extensively used over the last few decades to manufacture different components in fossil fired power generation plants, chemical and petro-chemical industries and in nuclear power plants [1–12]. Addition of Nb and V in modified 9Cr–1Mo steel provides superior creep resistance and high temperature strength under monotonic loading, compared to conventional 9Cr–1Mo steel [1–7]. Fine microalloy precipitates such as Nb(C, N) and V(C, N) retard the dislocation motion at high temperature and thereby enhance the creep strength of modified 9Cr–1Mo steel [1–7]. This alloy is presently being used as a stream-generator material in the Na-cooled fast breeder reactors [1–12]. Modified 9Cr–1Mo steel is the candidate material for use in the cladding of nuclear fuel, heat exchangers, pressure vessels and it is the primary tube material for nuclear reactors [1–7]. Modified 9Cr–1Mo steel shows higher thermal conductivity, lower thermal expansion coefficient and higher resistance to stress corrosion cracking in water-steam systems compared to austenitic stainless steels [1–7]. Selection of the modified 9Cr–1Mo steel for steam generator applications is also based on its good weldability and microstructural stability over prolonged exposure to high temperature service conditions [1–7].
n
Corresponding author.
http://dx.doi.org/10.1016/j.msea.2014.02.076 0921-5093 & 2014 Elsevier B.V. All rights reserved.
The components for high temperature systems are often subjected to repeated thermal stresses due to the temperature gradients arising from heating and cooling during start-ups and shut-downs. One of the major damage mechanisms for the nuclear reactors that operate at high temperature is fatigue by thermal gradients and by cyclic mechanical loading [1–21]. Hence, it is essential to evaluate the low-cycle fatigue properties of modified 9Cr–1Mo steel at ambient temperature, as well as at elevated (service) temperature, for designing the nuclear power plant components [1–18]. Several attempts have been made to characterize the LCF and creep–fatigue interaction behavior of modified 9Cr–1Mo steel [1–7,16–18]. Under cyclic loading, modified 9Cr–1Mo steel is known to exhibit extensive softening as a result of microstructural instability, which deteriorates its service life [1–7]. Considering the importance of cyclic softening phenomenon during elevated temperature service condition of modified 9Cr– 1Mo steel, the aim of the present study is to carry out a detailed investigation on low-cycle fatigue behavior of this steel. Both from phenomenological and operational points of view the prime interest is to understand the effect of temperature on low-cycle fatigue behavior and, therefore, three different test temperatures (25 1C, 500 1C, and 600 1C) were selected for the present study. The response of the indigenously developed modified 9Cr–1Mo steel under monotonic and cyclic loading over a range of test temperatures has been explained in view of the microstructural changes and fractographic analysis.
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Table 1 Chemical composition of modified 9Cr–1Mo steel as examined. Element
C
Si
Mn
P
S
Cr
Mo
Ni
Al
N
Nb
V
Fe
(wt%)
0.13
0.29
0.42
0.03
0.0034
9.03
0.94
0.07
0.017
0.03
0.073
0.185
88.80
2. Experimental details 2.1. Material The modified 9Cr–1Mo steel was obtained from BHEL R&D, India, in the form of a pipe having 500 mm length 120 mm OD 80 mm ID (OD: outer diameter, ID: inner diameter). Chemical composition of the steel showed the presence of V and Nb beside Cr and Mo, Table 1. The as-received material was in the normalized (1070 1C for 2 h) and tempered (740 1C for 2 h) conditions.
2.2. Tensile and fatigue testing Tensile and fatigue test specimens were machined from the longitudinal direction of the pipe section as per ASTM E8M standard and ASTM E606 standard, respectively [22,23]. Specimen surfaces were ground and polished and the tests were conducted at three different temperatures (25 1C, 500 1C and 600 1C). Calibrated thermo-couples were attached to the specimens within gauge length to monitor the test temperature at an accuracy of 73 1C. A 100 kN servo-electric universal testing machine (Instron 8860 model) having a three zone split furnace attached to it was used for performing all the tensile and fatigue tests. Displacement rates of 0.008 mm/s and 0.02 mm/s were applied for tensile testing at ambient and elevated temperatures, respectively. An extensometer of 12.5 mm length was mounted on the specimen gage length for strain measurement. Axial strain controlled low-cycle fatigue tests were performed at a constant strain-rate of 1 10 3 s 1 by varying the strain amplitude between 70.7% and 71.2%. About 200 data points were recorded per cycle for constructing the stress–strain hysteresis loop. Tests were conducted until a 20% load drop was noticed over the peak tensile load.
2.3. Microstructural evaluation Thin slices ( 1 mm) were cut from the gauge section of the fatigue tested specimens, polished down to 0.1 mm SiC emery paper and finally electro-polished in 10% perchloric acid in acetic acid solution. Thin foils ( 50 nm thick) were examined by JEOL 2000FX transmission electron microscope, operated at 200 kV.
3. Results and discussion 3.1. Microstructure of as-received steel The as-received microstructure of modified 9Cr–1Mo steel in normalized and tempered conditions showed typical tempered martensite structure, Fig. 1. As expected the structure consisted of prior-austenite grain boundaries, packet-, block- and lathboundaries and different types of carbide particles. The coarse particles were distributed along the lath-boundaries and prioraustenite grain boundaries (generally M23C6 type of precipitates) and the fine particles (MC type of precipitates) distributed uniformly within the lath structure, Fig. 1.
Fig. 1. Microstructure of modified 9Cr–1Mo steel in as-received (normalized and tempered) condition.
Table 2 Tensile properties of modified 9Cr–1Mo steel. Temperature (1C)
YS (MPa)
UTS (MPa)
% EL
% RA
UTS/YS ratio
25 500 600
650 504 365
790 579 399
27 15 19
69 75 90
1.21 1.14 1.09
3.2. Cyclic stress response as the function of temperature and strain amplitude The tensile test results of the investigated steel at three different test temperatures are presented in Table 2. An increase in test temperature resulted in a decrease in yield strength (YS) and ultimate tensile strength (UTS) and an increase in reduction in area (% RA), Table 2. The total elongation, however, was much higher in the sample tested at room temperature, compared to those tested at high temperatures, Table 2. The UTS:YS ratio decreased with an increase in test temperature, indicating a drop in the strain-hardening ability of the steel with respect to temperature, Table 2. At a given strain amplitude the steel exhibited rapid initial softening followed by gradual softening till the onset of final load drop, accompanied by the initiation and propagation of the fatigue crack, Fig. 2. The same trend was noticed at all the test temperatures and at different strain amplitudes, Fig. 2. The cyclic stress response of the investigated steel was found to decrease with the increase in test temperature, irrespective to the strain amplitude, Fig. 2(a). The half-life cycle stress–strain curve at any temperature can be represented by the following power law relationship [5]: 0 Δεp n Δs ¼ K0 ð1Þ 2 2 where Κ 0 is the cyclic strength coefficient, n0 is the cyclic strainhardening exponent, Δs is the amplitude of stress, and Δεp is the amplitude of plastic strain. The K 0 and n0 values obtained at different test temperatures are listed in Table 3.
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600 700 500
500
Stress, MPa
Stress, MPa
600
400 Δε= Δε= Δε= Δε= Δε= Δε=
300 200 100
+ 0.7% + 0.75% + 0.8% + 1.0% + 1.1% + 1.2%
400
Δε = Δε = Δε = Δε = Δε = Δε =
300
200
0
+ 0.7% + 0.75% + 0.8% +1.0% +1.1% +1.2%
100 0
200
400
600
800
1000 1200 1400
0
100
200
Cycle Δε = + Δε = + Δε = + Δε = + Δε = + Δε = +
700 600
Stress, MPa
300
400
500
600
Cycle
500
0.7% 0.75% 0.8% 1.0% 1.1% 1.2%
400 300 200 100 0 0
100
200
300
400
500
600
700
Cycle Fig. 2. Cyclic stress response of the fatigue tested samples for different test temperatures: (a) 25 1C, (b) 500 1C and (c) 600 1C.
Table 3 Cyclic stress–strain parameters of modified 9Cr–1Mo steel. Temperature (1C)
K0
n0
Kn
nn
25 500 600
1188 642 770
0.14 0.08 0.19
759 472 318
0.09 0.08 0.09
The cyclic stress–strain curve is determined by connecting the tips of the stable hysteresis loops of the fatigue tested specimens for different strain amplitudes and temperatures, Fig. 3. Comparison of the monotonic stress–strain curve and cyclic stress–strain curve points to the cyclic softening of the material at all the three test temperatures, Fig. 3. The present finding agrees with the hypothesis that cyclic softening is expected in the material that shows TS: YSr1.2 (YS and TS obtained from tensile test) [5,24,25]. 3.3. Analysis of hysteresis loop The following equation can be used to calculate the cyclic hardening (or softening) factor (Hs) considering the stable and initial hysteresis loop [17]: sa so ð2Þ Hs ¼
sa
where sa is the stable hysteresis loop and so is the initial hysteresis loop. A positive and negative value of Hs stands for the hardening and softening, respectively. Hardening/softening
factor determined from the hysteresis loop at stable cycle for 71% strain amplitudes showed complete softening behavior, Fig. 3, which is expected from the earlier studies on modified 9Cr–1Mo steel [4,5]. The hardening factor presented in Fig. 4 reflects rapid softening at the initial stage of hardening and the rate of softening gradually decreased and become saturated with the increasing number of cycles. Cyclic stress–strain curve can describe the relationship between stable stress and strain amplitudes but it cannot describe the branching in the hysteresis loop. Masing [24] suggested that the branching in hysteresis loop can be geometrically identical or it can be obtained by magnifying the cyclic stress–strain curve by a factor of two. In order to verify the Masing hypothesis in the stable hysteresis loops, the loops obtained at different strain amplitudes were translated to a common origin. If the upper branches of the stable hysteresis loops form a common envelope, then the material can be considered to follow the Masing behavior [24]. The envelope curves of the investigated steel samples represent the non-Masing behavior, irrespective of the test temperature, Fig. 5. At 25 1C, the stable hysteresis loops displayed near Masing relationship separately at lower strain-amplitudes (70.7%, 70.75% and 71.0%) and higher strain-amplitudes (71.1% to 71.2%); however, the behavior was non-Masing type when all the stable loops are considered together, Fig. 5(a). In a similar way at 500 1C, near-Masing behavior was noticed at the following strain-amplitudes: (i) 70.7%, 70.75%, 71.0% and 71.2% and (ii) 70.8% and 71.1%, Fig. 5(b). The same at 600 1C was found at the following strain-amplitudes: (i) 70.8% and 71.0% and (ii) 71.1% and 71.2%, Fig. 5(c). Overall, non-Masing behavior was observed at each test
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600
Stress, MPa
400
600
± 0.7% ± 0.75% ± 0.8% ± 1.0% ± 1.1% ± 1.2%
±0.7% ±0.75% ±0.8% ±1.0% ±1.1% ±1.2%
400
Stress, MPa
800
200 0 -200
199
200 0 -200
-400
-400 Cyclic Monotonic
-600
-800 -1.5 -1.2 -0.9 -0.6 -0.3 0.0 0.3 0.6 0.9 1.2 1.5
Cyclic monotonic -600 -1.5 -1.2 -0.9 -0.6 -0.3 0.0 0.3 0.6 0.9 1.2 1.5
Strain, %
Strain, %
400 300
Stress, MPa
200
± ± ± ± ± ±
100 0 -100 -200 -300 -400 -1.5 -1.2 -0.9 -0.6 -0.3 0.0
0.3
0.6
0.9
1.2
1.5
Strain, % Fig. 3. Comparison between monotonic and cyclic stress–strain behavior at different test temperatures: (a) 25 1C, (b) 500 1C and (c) 600 1C.
[5,24]
0.05 0.00
Δεn ¼
Hardening factor
-0.05 -0.10 -0.15 -0.20 -0.25 -0.30 25 C 500 C 600 C Strain amplitude: +1%
-0.35 -0.40 -0.45 0
50
100
150
200
250
Number of cycle Fig. 4. Hardening factor calculated at three different test temperatures for 7 1% strain amplitude.
temperature when all the strain-amplitudes are considered together. The nature of cyclic deformation behavior depends on the material composition, processing history and deformation micro-mechanism operating at the test temperature [24].
n ð1=nn Þ Δs n Δs þ2 E 2K n
ð3Þ
where asterisk (n) in superscript indicates that the corresponding parameter is measured with respect to the new origin (On). Kn and nn values are obtained from fitting to the experimental data, Table 3. The relationships between the old and new coordinate systems, as shown in Fig. 6, are available in the literature [5,24]. The master curve parameters, determined by the present study, decreased with the increase in test temperature, Table 3. 3.5. Study of Bauschinger effect and hysteresis loop If a specimen is deformed plastically beyond the yield stress in one direction (in tension), and then after unloading to zero stress and immediate reloading in the opposite direction (in compression) leads to the lowering of the yield stress upon reloading [25]. The above phenomenon is known as the Bauschinger effect and it is associated with the formation of dislocation network in polycrystalline metals and the breakdown of those networks upon unloading [25]. The Bauschinger strain (β) is defined as the plastic strain corresponding to the stress reversal at 75% of the maximum tensile stress in forward direction and it can be expressed as [16]
3.4. Analysis of master curves
β ¼ RΔεp
The master curve can be obtained by matching the upper braches of the hysteresis loops for different strain-amplitudes, by adjusting the origin of each loop such that the curve can envelope the ascending (loading) branches of hysteresis loops. The equation for the master curve with the new origin (On), corresponding to the loop with the minimum proportional range, can be given by
where R is the function of cyclic strain-hardening exponent as 0 given by R ¼ ð0:875Þ n [16]. The Bauschinger strain followed a single linear relationship with plastic strain range considering all three test temperatures, as shown in Fig. 7, which is anomalous to the present finding as a linear relationship is expected in metals that follow Masing behavior [16,24].
ð4Þ
K. Guguloth et al. / Materials Science & Engineering A 604 (2014) 196–206
1400
1400
1200
1200
1000
1000
800 600
± 0.7% ± 0.75% ± 0.8% ± 1.0% ± 1.1% ± 1.2%
400 200 0
Stress, MPa
Stress, MPa
200
± 0.7% ± 0.75% ± 0.8% ± 1.0% ± 1.1% ± 1.2%
800 600 400 200 0
0.0 0.3 0.6 0.9 1.2 1.5 1.8 2.1 2.4 2.7 3.0
0.0 0.3 0.6 0.9 1.2 1.5 1.8 2.1 2.4 2.7 3.0
Strain, %
Strain, %
1400
± 0.7% ± 0.75% ± 0.8% ± 1.0% ± 1.1% ± 1.2%
1200
Stress, MPa
1000 800 600 400 200 0
0.0 0.3 0.6 0.9 1.2 1.5 1.8 2.1 2.4 2.7 3.0 Strain, %
Fig. 5. Analysis of Masing behavior of the stable hysteresis loops for different test temperatures: (a) 25 1C, (b) 500 1C and (c) 600 1C.
1400
Stress, MPa
1000 800 600
700
800 800
600
± 0.7% ± 0.75% ± 0.8% ± 1.0% ± 1.1% ± 1.2%
400
400 200
200 δσ0
0
0 0.0
600 500
0.3
0.6
0.9
1.2
1.5
1.8
2.1
2.4
2.7
-200
Stress, MPa
1200
1000 1000
600 400
400
± 0.7% ± 0.75% ± 0.8% ± 1.0% ± 1.1% ± 1.2%
300 200
200 δσ0 0
100 0 0.0
0.3
0.6
0.9
1.2
1.5
1.8
2.1
2.4
2.7
-200
0.0 0.3 0.6 0.9 1.2 1.5 1.8 2.1 2.4 2.7 3.0
0.0 0.3 0.6 0.9 1.2 1.5 1.8 2.1 2.4 2.7 3.0
Strain, %
Strain, %
± ± ± ± ± ± δσ0
Fig. 6. Construction of master curves from the stable hysteresis loops for different test temperatures: (a) 25 1C, (b) 500 1C and (c) 600 1C.
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Table 4 Values of the constants in Coffin–Manson relationships at different test temperatures.
Fig. 7. Variation in Bauschinger strain as the function of plastic strain range for modified 9Cr–1Mo steel at different test temperatures.
Fig. 8. Coffin–Manson relationship of modified 9Cr–1Mo steel for different test temperatures: (a) 25 1C, (b) 500 1C and (c) 600 1C.
3.6. Study of fatigue life The modified 9Cr–1Mo steel showed a decrease in the fatigue life with the increase in test temperature, Fig. 2. The effect of temperature on the fatigue life was more prominent at lower strain amplitudes (less than 7 1.0%), Fig. 2. The usual way of representing the low-cycle fatigue test results is to plot the variation in plastic strain amplitude, ðΔεp =2Þ against the number of cycles to failure, (2Nf). This type of behavior is known as the Coffin–Manson relationship [5,24] and is presented in Fig. 8. The fatigue lives at all the test temperatures, as represented by the straight line in Fig. 8, were found to obey a modified Manson– Coffin relation [24,25]. Δεt Δεe Δεp sf ¼ þ ¼ ð2N f Þb þ ε0f ð2N f Þc 2 2 2 E 0
ð5Þ
where ðΔεt =2Þ is the total strain amplitude, ðΔεe =2Þ is the elastic strain amplitude, εf' is the fatigue ductility coefficient, c is the fatigue ductility exponent, s0f is the fatigue strength coefficient, b is the fatigue strength exponent and E is the elastic modulus. Values of different constants determined from Eq. (5) are listed in Table 4. 3.7. Microscopic study on fatigue tested specimens The TEM micrograph of the as-received steel showed a tempered martensite lath structure with a high dislocation density inside the laths, Fig. 9(a) and (b). Three types of precipitates were observed: (i) coarse and elliptical (100–500 nm length) Cr23C6 precipitates were preferentially located along the lath-, packetand prior-austenite grain boundaries, Fig. 9(b) and (c); (ii) cuboidal or faceted Mo-rich M2C type of precipitates (also containing Nb
Temperature (1C)
ε0f
c
s0f /E
b
25 500 600
0.075 0.208 0.160
0.36 0.53 0.47
0.0035 0.0033 0.0025
0.082 0.058 0.099
and V) of 50–150 nm size range were found within the martensitic laths and as well as along the lath boundaries, Fig. 9(b); and (iii) fine microalloy precipitates (carbides and carbo-nitrides of Nb and V) of less than 20 nm size were uniformly distributed throughout the microstructure, Fig. 9(d). Nature of the precipitates was identified from the SAED and EDS analyses, Fig. 9(e) and (f). After low-cycle fatigue testing at 25 1C with 71% strain amplitude the lath structure becomes unstable and it was converted to dislocation cell or sub-grain structure, Fig. 10(a)–(c). Lath width was also found to increase from 0.3–0.5 mm to 1.0– 1.5 mm, Fig. 10(a) and (b). The dislocation structure was highly heterogeneous with the packets of dislocation debris and illdefined cells present in it, Fig. 10(a)–(c). Cell size varied between 0.8 and 1.7 mm, Fig. 10(b) and (c). The precipitates located along the lath boundaries imposed a pinning effect, which resulted in a slightly elongated cell structure, Fig. 10(a)–(c). Equiaxed dislocation cells were also present, Fig. 10(b) and (c). Very fine microalloy precipitates were found to interact with the dislocations in dislocation-rich regions, Fig. 10(d), thereby retarding the dislocation motion. The microstructural features observed in the fatigue tested samples were in line with the earlier findings [2,7,9,16]. It is known that the dislocations rearrange themselves during cyclic loading into a lower energy configuration comprising of dislocation cells, surrounded by dislocation walls, by cross-slip [1–4]. Average dislocation density was measured to be (5 7 1.8) 104/ m2 in the as-received tube, which comes down to (1.5 70.7) 104/m2 after 261 stable cycles at room temperature. Increase in the test temperature to 500 1C resulted in the coarsening of dislocation cell structure, Fig. 11. Cr23C6 precipitates tend to coarsen and the adjacent precipitates coagulate and become more rounded in shape, Fig. 11(a). The dislocation density decreased further through the continuous interaction and annihilation of the dislocations of opposite signs. As per the earlier studies, significant difference in microstructural features are not expected below 550 1C, as dislocation climb cannot be fully operative [6–8]. TEM micrographs depicted the extensive interactions between the fine microalloy precipitates with the dislocations in the interlath regions, Fig. 11(b)–(d). Pronounced softening was detected at 600 1C due to the significant decrease in dislocation density to (0.170.003) 104/m2 and the density of fine, microalloy precipitates, Fig. 12. Cyclic softening is accompanied by the formation and growth of sub-grain structures due to the recovery, Fig. 12(a). Both M23C6 and MC type of precipitates tend to coarsen, Fig. 12(b) and (c), leading to the significant softening of the steel. The fine precipitates might also have dissolved as the propagating dislocations cut down the precipitates into a size smaller than the critical nucleus size [2]. Total cyclic stress amplitude (s) can be expressed in terms of internal stress (si) and effective stress (se), such as [9,10]
s ¼ se ðT; ε_ Þ þ si
ð6Þ
where the effective stress, se, is the thermally activated component and it depends on the short-range obstacles such as dislocation–precipitation interaction. Internal stress, si, is influenced by the dislocation–dislocation interaction and the interaction between the dislocations and the lath boundaries [10]. Now the
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Fig. 9. TEM micrographs of the as-received steel: (a) heavily dislocated and elongated martensitic lath structure; (b) elliptical shaped precipitates (arrowed) along with the dislocation networks; (c) coarse Cr23C6 precipitate (arrowed) and (d) fine NbC and VC precipitates (arrowed) interacting with the dislocations within the laths; (e) SAED pattern analysis and (f) EDS analysis identifying the precipitate in (b) within the circle as (Mo, Nb)C and the precipitate within the rectangular block in (b) contains Mo and V.
cyclic yield stress (sy) and the maximum cyclic stress (smax) can be represented by
sy ¼ se si
ð7Þ
smax ¼ se þ si
ð8Þ
Both se and si decreased with the increase in test temperature, accompanied by the decrease in precipitate density, dislocation density and the coarsening (and disappearance) of martensitic lath structure. As a result smax decreased with temperature (from 1100 to 1200 MPa at 25 1C to 550–600 MPa at 600 1C), Fig. 5. Decrease in se has a dominating effect over the decrease in si and hence sy also decreases with the increase in test temperature, Fig. 5. The cyclic strain hardening coefficients decreased with the increase in test temperature, Table 3. This is due to the continuous decrease in dislocation density and precipitate density, which reduced the dislocation–dislocation and dislocation–precipitate interactions. At 500 1C, dislocation density decreased and the fine MC type of precipitates continued to pin the existing dislocations.
As a result the number of mobile dislocations could have reduced significantly, which resulted in the pronounced drop in the cyclic strain hardening coefficients, Table 3. As the precipitates coarsened at 600 1C, the few dislocations present at that temperature become mobile, increasing the cyclic hardening coefficients slightly, Table 3. Increase in cyclic hardening coefficients at 600 1C can also be due to the occurrence of dynamic strain ageing, which is expected to take place in modified 9Cr–1Mo steel during fatigue testing at elevated temperatures [1,4,5]. 3.8. Fractographic studies on the fatigue tested specimens Fractographic study has been carried out on the fatigue tested specimens using scanning electron microscope (SEM). The fracture surfaces consisted of beach marks (Fig. 13(a) and (b)), which are the classical features of fatigue failure [1,2,5,11]. The failure mode was transgranular fracture, which comprised quasicleavage fracture of the martensitic laths, whilst plastic deformation (i.e. micro-void nucleation and growth) occurred along the lathboundaries, Fig. 13(b). Extensive branching of transgranular cracks
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Fig. 10. TEM micrographs of the sample fatigue tested at 25 1C: (a) coarsened (martensitic) lath structure having an elliptical shape precipitate along the lath boundary; (b and c) dislocation cell structure along with the coarse precipitates and (d) interaction between dislocations and the fine microalloy precipitates. Coarse and fineprecipitates are indicated by arrows.
Fig. 11. TEM micrographs of the sample fatigue tested at 500 1C: (a) coarsened dislocation cell structure having low dislocation density, along with precipitates (arrowed). Red arrow showing the coagulation of two precipitates; (b) bright field and (c) dark field images showing the uniform distribution of fine precipitates and (d) the corresponding SAED analysis identifying the precipitates to be V(C, N). (For interpretation of references to color in this figure legend, the reader is referred to the web version of this article.)
was noticed, especially at the higher strain amplitudes, and it was associated with the coarse-particles and inclusions. Fig. 13(c) shows such a crack branching and the presence of a coarse-precipitate
within the crack. During fatigue cycling of modified 9Cr–1Mo steel, micro-voids nucleate at the particle–matrix interface and grow in size under the stress-field of advancing crack [2,5,11].
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Fig. 12. TEM micrographs of the sample fatigue tested at 600 1C: (a) formation of sub-grain structure; (b and c) coarsened precipitates within the sub-grains along with few dislocations.
As the coarse-particles are primarily located at the lath boundaries, void initiation occurs at those locations. Eventually the voids grow and the stress build-up at the crack tip reaches sufficiently high values such that the inter-particle ligaments fail by quasi-cleavage fracture, advancing the crack. Formation of micro-voids around the cuboidal Mo-rich precipitates is evident from Fig. 13(d) and (e) and the corresponding EDS analysis from one such precipitates is shown in Fig. 13(f). Oxidation is expected to impose a detrimental effect on the initiation and propagation of fatigue crack at elevated temperatures [4,5,11]. Fig. 13(g) presents such an example, where oxide formed within the fatigue crack, propagating inwards from the surface of the sample tested at 600 1C. Rapid drop in the fatigue life at 600 1C can therefore be attributed to the oxidation phenomenon.
4. Conclusions Detailed investigation on the macroscopic behavior of modified 9Cr–1Mo steel under monotonic and cyclic loading (hysteresis loop) at three different test temperatures (25 1C, 500 1C and 600 1C) and the corresponding microstructural and fractographic studies lead to the following major conclusions: 1. Modified 9Cr–1Mo steel, having tempered martensitic structure containing coarse-precipitates (such as Cr23C6) and fine-precipitates (such as Mo2C, NbC and VC), showed the continuous softening to failure under cyclic loading both at ambient and at elevated temperatures. The degree of softening increased with the increase in test temperature.
2. Increase in test temperature also led to the continuous decrease in dislocation density by dislocation annihilation, coarsening of martensitic laths and the transformation of lath structure into dislocation cell and sub-grain structure. All these factors contributed to the cyclic softening of the steel. 3. Fine MC type of precipitates (M: V or Nb metal) interacted with the dislocations and retarded the dislocation motion at 25 1C and 500 1C test temperatures. Coarsening and dissolution of precipitates at 600 1C significantly decreased the precipitate pinning effect on the dislocations, leading to cyclic softening. 4. Strain-hardening coefficients decreased initially with the increase in test temperature from 25 1C and 500 1C, accompanied by the microstructural softening. Subsequent rise in the coefficients (at 600 1C) can be due to either the increase in mobile dislocation density (due to precipitate coarsening) or the dynamic strain ageing effect. 5. The investigated steel showed non-Masing behavior at ambient temperature. At elevated temperature, stable hysteresis loop shows partial Masing behavior due to the thermal activation. 6. In modified 9Cr–1Mo steel, the Bauschinger strain increased linearly with cyclic plastic strain. 7. At elevated temperatures, localized oxidation can be a major factor that contributes to cyclic fatigue damage. The formation of surface film oxides and fully reversed cyclic straining contributes to local increase in crack length and accelerates the crack propagation rate, leading to cyclic softening and premature softening. 8. Fractographic study on fatigue tested specimens showed the occurrence of quasi-cleavage fracture of martensitic laths and
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Fig. 13. SEM fractographs of fatigue tested specimens: (a and b) beach marks on the fracture surface; (c) quasi-cleavage facets and MnS inclusion situated within the secondary crack; (d and e) micro-voids forming around the cuboidal precipitates and the EDS analysis showing the precipitates to be Mo-rich (f); oxidation within the fatigue cracks (arrowed) propagating from the surface (g).
micro-void formation along the lath boundaries. Micro-voids were found to form around the precipitates, such as the cuboidal Mo2C particles. References [1] H.W. Zhou, Y.Z. He, H. Zhnag, Y.W. Cen, Int. J. Fatigue 47 (2013) 83–89. [2] V. Shankar, M. Valsan, K.B.S. Rao, R. Kannan, S.L. Manna, S.D. Pathak, Mater. Sci. Eng. A 437 (2006) 413–422.
[3] B. Fournier, M. Sauzay, A. Renault, F. Barcelo, A. Pineau, J. Nucl. Mater. 386–388 (2009) 71–74. [4] A. Nagesha, R. Kannan, G.V.S. Sastry, R. Sandhya, V. Singh, K. Bhanu Shankara Rao, Mater. Sci. Eng. A 554 (2012) 95–104. [5] A. Nagesha, M. Valsan, R. Kannan, K. Bhanu Sankara Rao, S.L. Mannan, Int. J. Fatigue 24 (2002) 1285–1293. [6] D.W. Kim, W.S. Ryu, Proc. Eng. 10 (2011) 376–382. [7] V. Shankar, V. Bauer, R. Sandhya, M.D. Mathew, H.J. Christ, J. Nucl. Mater. 420 (2012) 23–30. [8] M.F. Giordana, P.F. Giroux, I. Alvarez-Armas, M. Sauzay, A. Armas, T. Kruml, Mater. Sci. Eng. A 550 (2012) 103–111.
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[9] B. Fournier, M. Sauzay, A. Pineau, Int. J. Plast. 27 (2011) 1803–1816. [10] G.V.P. Reddy, R. Sandhya, K. Laha, M.D. Mathew, Mater. Sci. Eng. A 527 (2010) 6449–6453. [11] V. Chawsal, G. Sashikala, S.K. Ray, S.L. Mannan, B. Raj, Mater. Sci. Eng. A 395 (2005) 251–264. [12] P.J. Ennis, A.C. Filemonowicz, Sadhana 28 (2003) 709–730. [13] R.P. Skelton, H.J. Maier, H.J. Christ, Mater. Sci. Eng. A 238 (1997) 377–390. [14] J. Polak, M. Petrenec, S. Degallaix, Scr. Mater. 47 (2002) 731–736. [15] J. Polak, F. Fardoun, S. Degallaix, Mater. Sci. Eng. A 297 (2001) 144–153. [16] A. Plumtree, H.A. Abdel-Raouf, Int. J. Fatigue 23 (2001) 799–805. [17] H.X. Shang, H.J. Ding, Eng. Fract. Mech. 54 (1996) 1–9. [18] G. Ebi, A.J. McEvily, Fatigue Eng. Mater. Struct. 7 (1984) 299–314. [19] T.S. Srivatsan, S. Sriram, C. Daniels, Eng. Fract. Mech. 56 (1997) 531–550.
[20] G. Bernhart, G. Moulinier, O. Brucelle, D. Delanges, Int. J. Fatigue 21 (1999) 179–186. [21] A.F. Armas, C. Petersen, R. Schmitt, M. Avalos, I. Alvarez-Armas, J. Nucl. Mater. 307–311 (2002) 509–513. [22] Standard Practice for Tensile Testing of Metallic Materials, ASTM International, Designation: E08-M, Annual Book of ASTM Standards, vol. 03.01, 2009. [23] Standard Practice for Strain-Controlled Fatigue Testing, ASTM International, Designation: E606-04, Annual Book of ASTM Standards, vol. 03.01, 2009. [24] Fernand Ellyin, Fatigue Damage, Crack Growth and Life Prediction, first edition, Chapman & Hall, London, 1997. [25] G.E. Dieter, Mechanical Metallurgy, McGraw-Hill Book Company, London, UK, 1988.
Contents lists available at ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Low-cyclic fatigue behavior of modified 9Cr–1Mo steel at elevated temperature Krishna Guguloth a,n, S. Sivaprasad a, D. Chakrabarti b, S. Tarafder a a b
CSIR-National Metallurgical laboratory, Material Science and Technology Division, Jamshedpur 831007, India Department of Metallurgical and Materials Engineering, Indian Institute of Technology, Kharagpur 721302, India
art ic l e i nf o
a b s t r a c t
Article history: Received 7 January 2014 Received in revised form 18 February 2014 Accepted 21 February 2014 Available online 2 March 2014
The low-cycle fatigue behavior of indigenously developed modified 9Cr–1Mo steel has been evaluated using a constant strain rate (1 10 3 s 1) at ambient temperature (25 1C) and at elevated temperatures (500–600 1C) over the strain amplitudes varying between 70.7% and 71.2%. Cyclic stress response showed a gradual softening regime that ended in a stress plateau until complete failure of the specimens. The estimated fatigue life decreased with the increase in test temperature. The effect of temperature on fatigue life was more pronounced at lower strain amplitudes. The cyclic deformation behavior at different temperatures has been analyzed from hysteresis loop and also in view of the changes taking place in dislocation structure and dislocation–precipitation interaction. Evaluation of low-cycle fatigue properties of modified 9Cr–1Mo steel over a range of test temperature can help in designing components for in-core applications in fast breeder reactors and in super heaters for nuclear power plants. & 2014 Elsevier B.V. All rights reserved.
Keywords: Low-cycle fatigue Modified 9Cr–1Mo steel Hysteresis loop Masing behavior Dislocation Precipitation
1. Introduction Martensitic steels with 9–12 wt% Cr has been extensively used over the last few decades to manufacture different components in fossil fired power generation plants, chemical and petro-chemical industries and in nuclear power plants [1–12]. Addition of Nb and V in modified 9Cr–1Mo steel provides superior creep resistance and high temperature strength under monotonic loading, compared to conventional 9Cr–1Mo steel [1–7]. Fine microalloy precipitates such as Nb(C, N) and V(C, N) retard the dislocation motion at high temperature and thereby enhance the creep strength of modified 9Cr–1Mo steel [1–7]. This alloy is presently being used as a stream-generator material in the Na-cooled fast breeder reactors [1–12]. Modified 9Cr–1Mo steel is the candidate material for use in the cladding of nuclear fuel, heat exchangers, pressure vessels and it is the primary tube material for nuclear reactors [1–7]. Modified 9Cr–1Mo steel shows higher thermal conductivity, lower thermal expansion coefficient and higher resistance to stress corrosion cracking in water-steam systems compared to austenitic stainless steels [1–7]. Selection of the modified 9Cr–1Mo steel for steam generator applications is also based on its good weldability and microstructural stability over prolonged exposure to high temperature service conditions [1–7].
n
Corresponding author.
http://dx.doi.org/10.1016/j.msea.2014.02.076 0921-5093 & 2014 Elsevier B.V. All rights reserved.
The components for high temperature systems are often subjected to repeated thermal stresses due to the temperature gradients arising from heating and cooling during start-ups and shut-downs. One of the major damage mechanisms for the nuclear reactors that operate at high temperature is fatigue by thermal gradients and by cyclic mechanical loading [1–21]. Hence, it is essential to evaluate the low-cycle fatigue properties of modified 9Cr–1Mo steel at ambient temperature, as well as at elevated (service) temperature, for designing the nuclear power plant components [1–18]. Several attempts have been made to characterize the LCF and creep–fatigue interaction behavior of modified 9Cr–1Mo steel [1–7,16–18]. Under cyclic loading, modified 9Cr–1Mo steel is known to exhibit extensive softening as a result of microstructural instability, which deteriorates its service life [1–7]. Considering the importance of cyclic softening phenomenon during elevated temperature service condition of modified 9Cr– 1Mo steel, the aim of the present study is to carry out a detailed investigation on low-cycle fatigue behavior of this steel. Both from phenomenological and operational points of view the prime interest is to understand the effect of temperature on low-cycle fatigue behavior and, therefore, three different test temperatures (25 1C, 500 1C, and 600 1C) were selected for the present study. The response of the indigenously developed modified 9Cr–1Mo steel under monotonic and cyclic loading over a range of test temperatures has been explained in view of the microstructural changes and fractographic analysis.
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Table 1 Chemical composition of modified 9Cr–1Mo steel as examined. Element
C
Si
Mn
P
S
Cr
Mo
Ni
Al
N
Nb
V
Fe
(wt%)
0.13
0.29
0.42
0.03
0.0034
9.03
0.94
0.07
0.017
0.03
0.073
0.185
88.80
2. Experimental details 2.1. Material The modified 9Cr–1Mo steel was obtained from BHEL R&D, India, in the form of a pipe having 500 mm length 120 mm OD 80 mm ID (OD: outer diameter, ID: inner diameter). Chemical composition of the steel showed the presence of V and Nb beside Cr and Mo, Table 1. The as-received material was in the normalized (1070 1C for 2 h) and tempered (740 1C for 2 h) conditions.
2.2. Tensile and fatigue testing Tensile and fatigue test specimens were machined from the longitudinal direction of the pipe section as per ASTM E8M standard and ASTM E606 standard, respectively [22,23]. Specimen surfaces were ground and polished and the tests were conducted at three different temperatures (25 1C, 500 1C and 600 1C). Calibrated thermo-couples were attached to the specimens within gauge length to monitor the test temperature at an accuracy of 73 1C. A 100 kN servo-electric universal testing machine (Instron 8860 model) having a three zone split furnace attached to it was used for performing all the tensile and fatigue tests. Displacement rates of 0.008 mm/s and 0.02 mm/s were applied for tensile testing at ambient and elevated temperatures, respectively. An extensometer of 12.5 mm length was mounted on the specimen gage length for strain measurement. Axial strain controlled low-cycle fatigue tests were performed at a constant strain-rate of 1 10 3 s 1 by varying the strain amplitude between 70.7% and 71.2%. About 200 data points were recorded per cycle for constructing the stress–strain hysteresis loop. Tests were conducted until a 20% load drop was noticed over the peak tensile load.
2.3. Microstructural evaluation Thin slices ( 1 mm) were cut from the gauge section of the fatigue tested specimens, polished down to 0.1 mm SiC emery paper and finally electro-polished in 10% perchloric acid in acetic acid solution. Thin foils ( 50 nm thick) were examined by JEOL 2000FX transmission electron microscope, operated at 200 kV.
3. Results and discussion 3.1. Microstructure of as-received steel The as-received microstructure of modified 9Cr–1Mo steel in normalized and tempered conditions showed typical tempered martensite structure, Fig. 1. As expected the structure consisted of prior-austenite grain boundaries, packet-, block- and lathboundaries and different types of carbide particles. The coarse particles were distributed along the lath-boundaries and prioraustenite grain boundaries (generally M23C6 type of precipitates) and the fine particles (MC type of precipitates) distributed uniformly within the lath structure, Fig. 1.
Fig. 1. Microstructure of modified 9Cr–1Mo steel in as-received (normalized and tempered) condition.
Table 2 Tensile properties of modified 9Cr–1Mo steel. Temperature (1C)
YS (MPa)
UTS (MPa)
% EL
% RA
UTS/YS ratio
25 500 600
650 504 365
790 579 399
27 15 19
69 75 90
1.21 1.14 1.09
3.2. Cyclic stress response as the function of temperature and strain amplitude The tensile test results of the investigated steel at three different test temperatures are presented in Table 2. An increase in test temperature resulted in a decrease in yield strength (YS) and ultimate tensile strength (UTS) and an increase in reduction in area (% RA), Table 2. The total elongation, however, was much higher in the sample tested at room temperature, compared to those tested at high temperatures, Table 2. The UTS:YS ratio decreased with an increase in test temperature, indicating a drop in the strain-hardening ability of the steel with respect to temperature, Table 2. At a given strain amplitude the steel exhibited rapid initial softening followed by gradual softening till the onset of final load drop, accompanied by the initiation and propagation of the fatigue crack, Fig. 2. The same trend was noticed at all the test temperatures and at different strain amplitudes, Fig. 2. The cyclic stress response of the investigated steel was found to decrease with the increase in test temperature, irrespective to the strain amplitude, Fig. 2(a). The half-life cycle stress–strain curve at any temperature can be represented by the following power law relationship [5]: 0 Δεp n Δs ¼ K0 ð1Þ 2 2 where Κ 0 is the cyclic strength coefficient, n0 is the cyclic strainhardening exponent, Δs is the amplitude of stress, and Δεp is the amplitude of plastic strain. The K 0 and n0 values obtained at different test temperatures are listed in Table 3.
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600 700 500
500
Stress, MPa
Stress, MPa
600
400 Δε= Δε= Δε= Δε= Δε= Δε=
300 200 100
+ 0.7% + 0.75% + 0.8% + 1.0% + 1.1% + 1.2%
400
Δε = Δε = Δε = Δε = Δε = Δε =
300
200
0
+ 0.7% + 0.75% + 0.8% +1.0% +1.1% +1.2%
100 0
200
400
600
800
1000 1200 1400
0
100
200
Cycle Δε = + Δε = + Δε = + Δε = + Δε = + Δε = +
700 600
Stress, MPa
300
400
500
600
Cycle
500
0.7% 0.75% 0.8% 1.0% 1.1% 1.2%
400 300 200 100 0 0
100
200
300
400
500
600
700
Cycle Fig. 2. Cyclic stress response of the fatigue tested samples for different test temperatures: (a) 25 1C, (b) 500 1C and (c) 600 1C.
Table 3 Cyclic stress–strain parameters of modified 9Cr–1Mo steel. Temperature (1C)
K0
n0
Kn
nn
25 500 600
1188 642 770
0.14 0.08 0.19
759 472 318
0.09 0.08 0.09
The cyclic stress–strain curve is determined by connecting the tips of the stable hysteresis loops of the fatigue tested specimens for different strain amplitudes and temperatures, Fig. 3. Comparison of the monotonic stress–strain curve and cyclic stress–strain curve points to the cyclic softening of the material at all the three test temperatures, Fig. 3. The present finding agrees with the hypothesis that cyclic softening is expected in the material that shows TS: YSr1.2 (YS and TS obtained from tensile test) [5,24,25]. 3.3. Analysis of hysteresis loop The following equation can be used to calculate the cyclic hardening (or softening) factor (Hs) considering the stable and initial hysteresis loop [17]: sa so ð2Þ Hs ¼
sa
where sa is the stable hysteresis loop and so is the initial hysteresis loop. A positive and negative value of Hs stands for the hardening and softening, respectively. Hardening/softening
factor determined from the hysteresis loop at stable cycle for 71% strain amplitudes showed complete softening behavior, Fig. 3, which is expected from the earlier studies on modified 9Cr–1Mo steel [4,5]. The hardening factor presented in Fig. 4 reflects rapid softening at the initial stage of hardening and the rate of softening gradually decreased and become saturated with the increasing number of cycles. Cyclic stress–strain curve can describe the relationship between stable stress and strain amplitudes but it cannot describe the branching in the hysteresis loop. Masing [24] suggested that the branching in hysteresis loop can be geometrically identical or it can be obtained by magnifying the cyclic stress–strain curve by a factor of two. In order to verify the Masing hypothesis in the stable hysteresis loops, the loops obtained at different strain amplitudes were translated to a common origin. If the upper branches of the stable hysteresis loops form a common envelope, then the material can be considered to follow the Masing behavior [24]. The envelope curves of the investigated steel samples represent the non-Masing behavior, irrespective of the test temperature, Fig. 5. At 25 1C, the stable hysteresis loops displayed near Masing relationship separately at lower strain-amplitudes (70.7%, 70.75% and 71.0%) and higher strain-amplitudes (71.1% to 71.2%); however, the behavior was non-Masing type when all the stable loops are considered together, Fig. 5(a). In a similar way at 500 1C, near-Masing behavior was noticed at the following strain-amplitudes: (i) 70.7%, 70.75%, 71.0% and 71.2% and (ii) 70.8% and 71.1%, Fig. 5(b). The same at 600 1C was found at the following strain-amplitudes: (i) 70.8% and 71.0% and (ii) 71.1% and 71.2%, Fig. 5(c). Overall, non-Masing behavior was observed at each test
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600
Stress, MPa
400
600
± 0.7% ± 0.75% ± 0.8% ± 1.0% ± 1.1% ± 1.2%
±0.7% ±0.75% ±0.8% ±1.0% ±1.1% ±1.2%
400
Stress, MPa
800
200 0 -200
199
200 0 -200
-400
-400 Cyclic Monotonic
-600
-800 -1.5 -1.2 -0.9 -0.6 -0.3 0.0 0.3 0.6 0.9 1.2 1.5
Cyclic monotonic -600 -1.5 -1.2 -0.9 -0.6 -0.3 0.0 0.3 0.6 0.9 1.2 1.5
Strain, %
Strain, %
400 300
Stress, MPa
200
± ± ± ± ± ±
100 0 -100 -200 -300 -400 -1.5 -1.2 -0.9 -0.6 -0.3 0.0
0.3
0.6
0.9
1.2
1.5
Strain, % Fig. 3. Comparison between monotonic and cyclic stress–strain behavior at different test temperatures: (a) 25 1C, (b) 500 1C and (c) 600 1C.
[5,24]
0.05 0.00
Δεn ¼
Hardening factor
-0.05 -0.10 -0.15 -0.20 -0.25 -0.30 25 C 500 C 600 C Strain amplitude: +1%
-0.35 -0.40 -0.45 0
50
100
150
200
250
Number of cycle Fig. 4. Hardening factor calculated at three different test temperatures for 7 1% strain amplitude.
temperature when all the strain-amplitudes are considered together. The nature of cyclic deformation behavior depends on the material composition, processing history and deformation micro-mechanism operating at the test temperature [24].
n ð1=nn Þ Δs n Δs þ2 E 2K n
ð3Þ
where asterisk (n) in superscript indicates that the corresponding parameter is measured with respect to the new origin (On). Kn and nn values are obtained from fitting to the experimental data, Table 3. The relationships between the old and new coordinate systems, as shown in Fig. 6, are available in the literature [5,24]. The master curve parameters, determined by the present study, decreased with the increase in test temperature, Table 3. 3.5. Study of Bauschinger effect and hysteresis loop If a specimen is deformed plastically beyond the yield stress in one direction (in tension), and then after unloading to zero stress and immediate reloading in the opposite direction (in compression) leads to the lowering of the yield stress upon reloading [25]. The above phenomenon is known as the Bauschinger effect and it is associated with the formation of dislocation network in polycrystalline metals and the breakdown of those networks upon unloading [25]. The Bauschinger strain (β) is defined as the plastic strain corresponding to the stress reversal at 75% of the maximum tensile stress in forward direction and it can be expressed as [16]
3.4. Analysis of master curves
β ¼ RΔεp
The master curve can be obtained by matching the upper braches of the hysteresis loops for different strain-amplitudes, by adjusting the origin of each loop such that the curve can envelope the ascending (loading) branches of hysteresis loops. The equation for the master curve with the new origin (On), corresponding to the loop with the minimum proportional range, can be given by
where R is the function of cyclic strain-hardening exponent as 0 given by R ¼ ð0:875Þ n [16]. The Bauschinger strain followed a single linear relationship with plastic strain range considering all three test temperatures, as shown in Fig. 7, which is anomalous to the present finding as a linear relationship is expected in metals that follow Masing behavior [16,24].
ð4Þ
K. Guguloth et al. / Materials Science & Engineering A 604 (2014) 196–206
1400
1400
1200
1200
1000
1000
800 600
± 0.7% ± 0.75% ± 0.8% ± 1.0% ± 1.1% ± 1.2%
400 200 0
Stress, MPa
Stress, MPa
200
± 0.7% ± 0.75% ± 0.8% ± 1.0% ± 1.1% ± 1.2%
800 600 400 200 0
0.0 0.3 0.6 0.9 1.2 1.5 1.8 2.1 2.4 2.7 3.0
0.0 0.3 0.6 0.9 1.2 1.5 1.8 2.1 2.4 2.7 3.0
Strain, %
Strain, %
1400
± 0.7% ± 0.75% ± 0.8% ± 1.0% ± 1.1% ± 1.2%
1200
Stress, MPa
1000 800 600 400 200 0
0.0 0.3 0.6 0.9 1.2 1.5 1.8 2.1 2.4 2.7 3.0 Strain, %
Fig. 5. Analysis of Masing behavior of the stable hysteresis loops for different test temperatures: (a) 25 1C, (b) 500 1C and (c) 600 1C.
1400
Stress, MPa
1000 800 600
700
800 800
600
± 0.7% ± 0.75% ± 0.8% ± 1.0% ± 1.1% ± 1.2%
400
400 200
200 δσ0
0
0 0.0
600 500
0.3
0.6
0.9
1.2
1.5
1.8
2.1
2.4
2.7
-200
Stress, MPa
1200
1000 1000
600 400
400
± 0.7% ± 0.75% ± 0.8% ± 1.0% ± 1.1% ± 1.2%
300 200
200 δσ0 0
100 0 0.0
0.3
0.6
0.9
1.2
1.5
1.8
2.1
2.4
2.7
-200
0.0 0.3 0.6 0.9 1.2 1.5 1.8 2.1 2.4 2.7 3.0
0.0 0.3 0.6 0.9 1.2 1.5 1.8 2.1 2.4 2.7 3.0
Strain, %
Strain, %
± ± ± ± ± ± δσ0
Fig. 6. Construction of master curves from the stable hysteresis loops for different test temperatures: (a) 25 1C, (b) 500 1C and (c) 600 1C.
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Table 4 Values of the constants in Coffin–Manson relationships at different test temperatures.
Fig. 7. Variation in Bauschinger strain as the function of plastic strain range for modified 9Cr–1Mo steel at different test temperatures.
Fig. 8. Coffin–Manson relationship of modified 9Cr–1Mo steel for different test temperatures: (a) 25 1C, (b) 500 1C and (c) 600 1C.
3.6. Study of fatigue life The modified 9Cr–1Mo steel showed a decrease in the fatigue life with the increase in test temperature, Fig. 2. The effect of temperature on the fatigue life was more prominent at lower strain amplitudes (less than 7 1.0%), Fig. 2. The usual way of representing the low-cycle fatigue test results is to plot the variation in plastic strain amplitude, ðΔεp =2Þ against the number of cycles to failure, (2Nf). This type of behavior is known as the Coffin–Manson relationship [5,24] and is presented in Fig. 8. The fatigue lives at all the test temperatures, as represented by the straight line in Fig. 8, were found to obey a modified Manson– Coffin relation [24,25]. Δεt Δεe Δεp sf ¼ þ ¼ ð2N f Þb þ ε0f ð2N f Þc 2 2 2 E 0
ð5Þ
where ðΔεt =2Þ is the total strain amplitude, ðΔεe =2Þ is the elastic strain amplitude, εf' is the fatigue ductility coefficient, c is the fatigue ductility exponent, s0f is the fatigue strength coefficient, b is the fatigue strength exponent and E is the elastic modulus. Values of different constants determined from Eq. (5) are listed in Table 4. 3.7. Microscopic study on fatigue tested specimens The TEM micrograph of the as-received steel showed a tempered martensite lath structure with a high dislocation density inside the laths, Fig. 9(a) and (b). Three types of precipitates were observed: (i) coarse and elliptical (100–500 nm length) Cr23C6 precipitates were preferentially located along the lath-, packetand prior-austenite grain boundaries, Fig. 9(b) and (c); (ii) cuboidal or faceted Mo-rich M2C type of precipitates (also containing Nb
Temperature (1C)
ε0f
c
s0f /E
b
25 500 600
0.075 0.208 0.160
0.36 0.53 0.47
0.0035 0.0033 0.0025
0.082 0.058 0.099
and V) of 50–150 nm size range were found within the martensitic laths and as well as along the lath boundaries, Fig. 9(b); and (iii) fine microalloy precipitates (carbides and carbo-nitrides of Nb and V) of less than 20 nm size were uniformly distributed throughout the microstructure, Fig. 9(d). Nature of the precipitates was identified from the SAED and EDS analyses, Fig. 9(e) and (f). After low-cycle fatigue testing at 25 1C with 71% strain amplitude the lath structure becomes unstable and it was converted to dislocation cell or sub-grain structure, Fig. 10(a)–(c). Lath width was also found to increase from 0.3–0.5 mm to 1.0– 1.5 mm, Fig. 10(a) and (b). The dislocation structure was highly heterogeneous with the packets of dislocation debris and illdefined cells present in it, Fig. 10(a)–(c). Cell size varied between 0.8 and 1.7 mm, Fig. 10(b) and (c). The precipitates located along the lath boundaries imposed a pinning effect, which resulted in a slightly elongated cell structure, Fig. 10(a)–(c). Equiaxed dislocation cells were also present, Fig. 10(b) and (c). Very fine microalloy precipitates were found to interact with the dislocations in dislocation-rich regions, Fig. 10(d), thereby retarding the dislocation motion. The microstructural features observed in the fatigue tested samples were in line with the earlier findings [2,7,9,16]. It is known that the dislocations rearrange themselves during cyclic loading into a lower energy configuration comprising of dislocation cells, surrounded by dislocation walls, by cross-slip [1–4]. Average dislocation density was measured to be (5 7 1.8) 104/ m2 in the as-received tube, which comes down to (1.5 70.7) 104/m2 after 261 stable cycles at room temperature. Increase in the test temperature to 500 1C resulted in the coarsening of dislocation cell structure, Fig. 11. Cr23C6 precipitates tend to coarsen and the adjacent precipitates coagulate and become more rounded in shape, Fig. 11(a). The dislocation density decreased further through the continuous interaction and annihilation of the dislocations of opposite signs. As per the earlier studies, significant difference in microstructural features are not expected below 550 1C, as dislocation climb cannot be fully operative [6–8]. TEM micrographs depicted the extensive interactions between the fine microalloy precipitates with the dislocations in the interlath regions, Fig. 11(b)–(d). Pronounced softening was detected at 600 1C due to the significant decrease in dislocation density to (0.170.003) 104/m2 and the density of fine, microalloy precipitates, Fig. 12. Cyclic softening is accompanied by the formation and growth of sub-grain structures due to the recovery, Fig. 12(a). Both M23C6 and MC type of precipitates tend to coarsen, Fig. 12(b) and (c), leading to the significant softening of the steel. The fine precipitates might also have dissolved as the propagating dislocations cut down the precipitates into a size smaller than the critical nucleus size [2]. Total cyclic stress amplitude (s) can be expressed in terms of internal stress (si) and effective stress (se), such as [9,10]
s ¼ se ðT; ε_ Þ þ si
ð6Þ
where the effective stress, se, is the thermally activated component and it depends on the short-range obstacles such as dislocation–precipitation interaction. Internal stress, si, is influenced by the dislocation–dislocation interaction and the interaction between the dislocations and the lath boundaries [10]. Now the
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Fig. 9. TEM micrographs of the as-received steel: (a) heavily dislocated and elongated martensitic lath structure; (b) elliptical shaped precipitates (arrowed) along with the dislocation networks; (c) coarse Cr23C6 precipitate (arrowed) and (d) fine NbC and VC precipitates (arrowed) interacting with the dislocations within the laths; (e) SAED pattern analysis and (f) EDS analysis identifying the precipitate in (b) within the circle as (Mo, Nb)C and the precipitate within the rectangular block in (b) contains Mo and V.
cyclic yield stress (sy) and the maximum cyclic stress (smax) can be represented by
sy ¼ se si
ð7Þ
smax ¼ se þ si
ð8Þ
Both se and si decreased with the increase in test temperature, accompanied by the decrease in precipitate density, dislocation density and the coarsening (and disappearance) of martensitic lath structure. As a result smax decreased with temperature (from 1100 to 1200 MPa at 25 1C to 550–600 MPa at 600 1C), Fig. 5. Decrease in se has a dominating effect over the decrease in si and hence sy also decreases with the increase in test temperature, Fig. 5. The cyclic strain hardening coefficients decreased with the increase in test temperature, Table 3. This is due to the continuous decrease in dislocation density and precipitate density, which reduced the dislocation–dislocation and dislocation–precipitate interactions. At 500 1C, dislocation density decreased and the fine MC type of precipitates continued to pin the existing dislocations.
As a result the number of mobile dislocations could have reduced significantly, which resulted in the pronounced drop in the cyclic strain hardening coefficients, Table 3. As the precipitates coarsened at 600 1C, the few dislocations present at that temperature become mobile, increasing the cyclic hardening coefficients slightly, Table 3. Increase in cyclic hardening coefficients at 600 1C can also be due to the occurrence of dynamic strain ageing, which is expected to take place in modified 9Cr–1Mo steel during fatigue testing at elevated temperatures [1,4,5]. 3.8. Fractographic studies on the fatigue tested specimens Fractographic study has been carried out on the fatigue tested specimens using scanning electron microscope (SEM). The fracture surfaces consisted of beach marks (Fig. 13(a) and (b)), which are the classical features of fatigue failure [1,2,5,11]. The failure mode was transgranular fracture, which comprised quasicleavage fracture of the martensitic laths, whilst plastic deformation (i.e. micro-void nucleation and growth) occurred along the lathboundaries, Fig. 13(b). Extensive branching of transgranular cracks
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Fig. 10. TEM micrographs of the sample fatigue tested at 25 1C: (a) coarsened (martensitic) lath structure having an elliptical shape precipitate along the lath boundary; (b and c) dislocation cell structure along with the coarse precipitates and (d) interaction between dislocations and the fine microalloy precipitates. Coarse and fineprecipitates are indicated by arrows.
Fig. 11. TEM micrographs of the sample fatigue tested at 500 1C: (a) coarsened dislocation cell structure having low dislocation density, along with precipitates (arrowed). Red arrow showing the coagulation of two precipitates; (b) bright field and (c) dark field images showing the uniform distribution of fine precipitates and (d) the corresponding SAED analysis identifying the precipitates to be V(C, N). (For interpretation of references to color in this figure legend, the reader is referred to the web version of this article.)
was noticed, especially at the higher strain amplitudes, and it was associated with the coarse-particles and inclusions. Fig. 13(c) shows such a crack branching and the presence of a coarse-precipitate
within the crack. During fatigue cycling of modified 9Cr–1Mo steel, micro-voids nucleate at the particle–matrix interface and grow in size under the stress-field of advancing crack [2,5,11].
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Fig. 12. TEM micrographs of the sample fatigue tested at 600 1C: (a) formation of sub-grain structure; (b and c) coarsened precipitates within the sub-grains along with few dislocations.
As the coarse-particles are primarily located at the lath boundaries, void initiation occurs at those locations. Eventually the voids grow and the stress build-up at the crack tip reaches sufficiently high values such that the inter-particle ligaments fail by quasi-cleavage fracture, advancing the crack. Formation of micro-voids around the cuboidal Mo-rich precipitates is evident from Fig. 13(d) and (e) and the corresponding EDS analysis from one such precipitates is shown in Fig. 13(f). Oxidation is expected to impose a detrimental effect on the initiation and propagation of fatigue crack at elevated temperatures [4,5,11]. Fig. 13(g) presents such an example, where oxide formed within the fatigue crack, propagating inwards from the surface of the sample tested at 600 1C. Rapid drop in the fatigue life at 600 1C can therefore be attributed to the oxidation phenomenon.
4. Conclusions Detailed investigation on the macroscopic behavior of modified 9Cr–1Mo steel under monotonic and cyclic loading (hysteresis loop) at three different test temperatures (25 1C, 500 1C and 600 1C) and the corresponding microstructural and fractographic studies lead to the following major conclusions: 1. Modified 9Cr–1Mo steel, having tempered martensitic structure containing coarse-precipitates (such as Cr23C6) and fine-precipitates (such as Mo2C, NbC and VC), showed the continuous softening to failure under cyclic loading both at ambient and at elevated temperatures. The degree of softening increased with the increase in test temperature.
2. Increase in test temperature also led to the continuous decrease in dislocation density by dislocation annihilation, coarsening of martensitic laths and the transformation of lath structure into dislocation cell and sub-grain structure. All these factors contributed to the cyclic softening of the steel. 3. Fine MC type of precipitates (M: V or Nb metal) interacted with the dislocations and retarded the dislocation motion at 25 1C and 500 1C test temperatures. Coarsening and dissolution of precipitates at 600 1C significantly decreased the precipitate pinning effect on the dislocations, leading to cyclic softening. 4. Strain-hardening coefficients decreased initially with the increase in test temperature from 25 1C and 500 1C, accompanied by the microstructural softening. Subsequent rise in the coefficients (at 600 1C) can be due to either the increase in mobile dislocation density (due to precipitate coarsening) or the dynamic strain ageing effect. 5. The investigated steel showed non-Masing behavior at ambient temperature. At elevated temperature, stable hysteresis loop shows partial Masing behavior due to the thermal activation. 6. In modified 9Cr–1Mo steel, the Bauschinger strain increased linearly with cyclic plastic strain. 7. At elevated temperatures, localized oxidation can be a major factor that contributes to cyclic fatigue damage. The formation of surface film oxides and fully reversed cyclic straining contributes to local increase in crack length and accelerates the crack propagation rate, leading to cyclic softening and premature softening. 8. Fractographic study on fatigue tested specimens showed the occurrence of quasi-cleavage fracture of martensitic laths and
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Fig. 13. SEM fractographs of fatigue tested specimens: (a and b) beach marks on the fracture surface; (c) quasi-cleavage facets and MnS inclusion situated within the secondary crack; (d and e) micro-voids forming around the cuboidal precipitates and the EDS analysis showing the precipitates to be Mo-rich (f); oxidation within the fatigue cracks (arrowed) propagating from the surface (g).
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