- Email: [email protected]
Creep–fatigue-oxidation interaction in Grade 91 steel weld joints for high temperature applications
Materials Science and Engineering A 528 (2011) 8428–8437
Contents lists available at ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Creep–fatigue-oxidation interaction in Grade 91 steel weld joints for high temperature applications Vani Shankar ∗ , R. Sandhya, M.D. Mathew Mechanical Metallurgy Division, Indira Gandhi Centre for Atomic Research, Kalpakkam 603102, Tamil Nadu, India
a r t i c l e
i n f o
Article history: Received 3 February 2011 Received in revised form 12 July 2011 Accepted 22 July 2011 Available online 29 July 2011 Keywords: Grade 91 steel weld joint Type IV cracking Low cycle fatigue Creep–fatigue interaction Oxidation
a b s t r a c t The weld joint consisting of a heterogeneous microstructure exhibits a lower fatigue life than that of the base metal. This is due to the presence of soft zone in the heat affected zone (HAZ). At high temperatures and under hold application, strain localization occurs in the soft intercritical HAZ (ICHAZ). Sub-surface creep cavity formation in the soft region and their linkage causes enhanced crack propagation and that translates into lower fatigue life of the weld joint at high temperatures. Occurrence of compression dwell sensitivity in the material is attributed to the presence of surface oxides. © 2011 Elsevier B.V. All rights reserved.
1. Introduction Modified version of plain 9Cr–1Mo steel in which Nb and V are optimized (Grade 91 steel) are being extensively used for super heater tubing, headers and piping of conventional as well as nuclear power plants with steam temperatures up to 866 K [1,2]. Compared to the austenitic stainless steels, the alloy has better resistance to thermal fatigue and stress corrosion cracking. The Grade 91 steel has better monotonic tensile, creep and fatigue properties at elevated temperatures compared to plain 9Cr–1Mo steel. Steam generators (SG) are complex large structures that are made in parts and integrated by welding techniques. Welds are weak links in any structure and many failures in high temperature components are reported to be weld-related [3–5]. During welding, the temperature sensitive microstructure of Grade 91 steel gets modified to various extents depending upon the temperature it is exposed to during the weld thermal cycles; highest temperature is reached by HAZ closest to the weld metal. Hence the narrow band of heat affected zone (HAZ) is a complex variation of microstructure with varied mechanical properties. Many premature failures at weld joints (WJ) have been identified to occur in a narrow fine grained or inter critical region of the HAZ in stronger 9–12 Cr steels such as Grade 91 ferritic steel [3–14] and they are termed as type IV failures.
∗ Corresponding author. Tel.: +91 44 274800x21147; fax: +91 44 27480075. E-mail address: [email protected] (V. Shankar). 0921-5093/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2011.07.046
In sodium cooled fast reactors (SFRs), the components operating at high temperatures are subjected to heating and cooling transients (sharp thermal gradients from surface to core of thick components) during startup and shutdown. This onsets thermal stresses in material if thermal strain (expansion/contraction) is totally/partly constrained. Repetition of these thermal transient results in strain-controlled low cycle fatigue of the components. In addition, steady state operation at elevated temperatures introduces creep, resulting in creep–fatigue interaction condition. In order to have more reliability in the performance of a component, it is essential to understand the deformation mechanisms operative under pure fatigue, pure creep and creep–fatigue interaction conditions. Creep experiments on weld joints of Grade 91 steel have shown that fracture usually occurs in the base metal in the short term tests at high stresses, whereas under long term creep conditions at low stresses Type IV failure mode appears [7–10]. In high chromium ferritic martensitic steels, significant reduction in rupture life as compared to the specimens without weld joint have been reported due to Type IV failure [6,9–14] especially at low stresses. Under creep at high stresses, failure has been reported to occur in the weld metal [15–18]. Whereas the creep performance of weld joints of Grade 91 steel is much better understood, the low cycle fatigue (LCF) and creep–fatigue interaction (CFI) behavior of the weld joints are not so well established. Hence the aim of the present paper is to study creep–fatigue interaction (CFI) behavior, identify the underlying deformation mechanism and correlate with the failure mode occurring in Grade 91 steel weld joints.
V. Shankar et al. / Materials Science and Engineering A 528 (2011) 8428–8437
8429
Table 1 Chemical composition (wt%) of base metal and weld metal. Element
C
Cr
Mo
Ni
V
Nb
N
S
P
Cu
Co
Base metal Weld metal
0.11 0.1
9.3 9
0.99 1.0
0.14 0.7
0.25 0.17
0.1 0.06
0.068 0.055
0.008 0.012
0.020 0.009
<0.05
0.016
2. Experimental Chemical composition of the 30 mm thick rolled plate used for current study and the weld metal is given in Table 1. Due to the high reliability required for nuclear components, the material has closely controlled composition, low residual element concentration and low inclusion content. The base material was given a normalizing treatment at 1313 K for 1 h followed by air cooling and subsequently tempering at 1033 K for 1 h followed by air cooling, before machining the fatigue specimens. Microstructural characterizations were carried out using optical microscope, scanning and transmission electron microscopes (TEM). Samples for the optical metallography were etched using Vilella’s reagent (1 g of picric acid + 5 ml conc. HCl + 100 ml ethyl alcohol). Samples for TEM were prepared by mechanically thinning the foils to 100 m thickness and finally electrolytic thinning at 20 V using stainless steel electrodes in a double jet electropolisher. The electrolyte used was 20% perchloric acid in methanol at 233 K. For fabricating LCF specimens containing weld joint, the plates were joined along the rolling direction by shielded metal arc welding (SMAW) process using voltage and current approximately 20 V and 100 A, respectively. A double-V configuration, with an included angle of 70◦ , a root face of 2 mm, and a root gap of 3.15 mm was used. Multiple passes were employed to fill the groove. An inter-pass temperature of 423 K was maintained during welding. Matching filler wire was used for welding. The weld pads were examined by radiography for their soundness. A nearly symmetrical weld profile was obtained after multipass welding (Fig. 1(a)). Specimens were taken from the centre of the welded plate (Fig. 1(b)). A schematic of specimen geometry containing the weld joint is shown in Fig. 1(c). The weld joint consists of three regions—base metal (BM), heat affected zone (HAZ) and weld metal (WM). The standard LCF specimen was selected in such a way that the WM was in the centre of the gauge length and the two HAZs were in the two sides of the WM and the rest of the sample was BM. Using metallographic and non destructive techniques, the double V weld profile, different zones of the HAZ and the amount of BM contained in the sample was first established. The maximum width of WM was around 10 mm with ∼3.5–4 mm HAZ and 3 mm BM on each side of WM. Bars of 110 mm length and 25 mm × 25 mm square cross section were cut from the welded plate. Post welding heat treatment (PWHT) was given to the bars containing weld joints at 1033 K for 3 h followed by air cooling. A nearly symmetrical weld profile was obtained (Fig. 1(b)). Low cycle fatigue (LCF) and creep–fatigue interaction (CFI) tests were conducted in air, under fully reversed, total axial strain control mode in accordance with ASTM specification E606 [19] in a
closed loop servo hydraulic testing system equipped with a radiant heating furnace. The temperature variation along the gauge length of the specimen did not exceed ±2 K. Continuous cycling tests were performed using a triangular waveform and hold time experiments were carried out using a trapezoidal waveform. LCF tests under continuous cycling condition were carried out at room temperature at strain amplitudes of ±0.6% and at 873 K at strain amplitude ±0.25%, ±0.4%, ±0.6% and ±1%. CFI tests were conducted by introducing hold of 1, 10 and 30 min at peak tension and 1 and 10 min in peak compression under ±0.6% strain amplitude. All the tests were carried out at a constant strain rate of 3 × 10−3 s−1 . A table summarizing the testing conditions is given as Table 2. After fatigue tests, the failed samples were pulled out to separate the fractured surface. Fractographic analysis was performed using SEM after which the longitudinal sections of the two separated parts were analyzed using an optical microscope equipped with an image analyzer. Microstructure near the fractured surface was studied for all the tested specimens. Microhardness was taken across the fusion line using 200 gmf load in a microhardness tester. Since failure usually occurred in one of the two sides of the weld joint (WJ), the microhardness measurements was performed in the side that had not cracked.
3. Results and discussion 3.1. Initial characterization of material The base material after normalizing and tempering treatment reveal a tempered martensitic microstructure with extensive precipitation of M23 C6 at the boundaries and MX type precipitates within the matrix [20]. Starting from weld metal towards base material, the microstructure is as follows: weld metal (Fig. 2(a)) with a tempered martensitic structure, tempered martensite structure with islands of ␦-ferrite (Fig. 2(b)), fine grained region (Fig. 2(c)), inter-critical structure (Fig. 2(d)) and unaffected base material (Fig. 2(e)) with a tempered martensitic structure. The total width of the HAZ is ∼3–4 mm. The intercritical region is approximately 2.5–3 mm away from the fusion line and close to BM. The region closest to the fusion boundary experiences a peak temperature above AC4 (the boundary between ␥ and ␥ + ␦ phase fields). At this high temperature, ␦ ferrite form along the austenite grain boundaries. The high alloy content of the steel resulting from dissolution of carbonitrides [21], coupled with the high cooling rate in the weld thermal cycle causes the retention of some ␦ ferrite on cooling. The resulting structure after cooling is therefore
Fig. 1. Weld profile obtained after multipass welding (a), details of weld joint specimens taken from welded plate (b) and schematic of specimen geometry containing the weld joint (c).
8430
V. Shankar et al. / Materials Science and Engineering A 528 (2011) 8428–8437
Table 2 Summary of results showing effect of temperature, strain amplitude and application of hold on LCF properties. Temperature (K)
Strain amplitude (%)
Type and duration of hold
1st cycle stress (MPa)
Peak stress at half-life (MPa)
300
0.6
0
572
520
1030
0 0 0
823
0.25 0.4 0.6 (BM) 0.6 1.0 0.6 0.6
328 375 456 392 410 354 389
262 291 315 300 296 271 307
3206 1007 580 475 209 374 240
0.25 0.4 0.6 (BM) 0.6 1.0 0.6 0.6 0.6 0.6
0 0 0
299 340 386 360 322 331 345 340 342
220 239 260 240 245 212 237 251 232
2250 1218 500 534 223 510 400 235 324
873
0 1 min TH 1 min CH
0 1 min TH 10 min TH 1 min CH 10 min CH
Fatigue life
Fig. 2. Resultant graded microstructure in the Grade 91 steel weld joint due to thermal exposure to different temperatures during weld thermal cycles; optical micrographs of different regions of weld joint such as weld metal (a), coarse grain with ␦ ferrite (b), fine grained region (c), intercritical region (d) and base metal (e).
coarse prior austenite grain martensite with ␦ ferrite along the prior austenite grain boundaries. Further away from the fusion boundary, where the peak temperature was well above AC3 but below AC4 , the carbides that impede austenite grain growth get dissolved and hence coarse grain austenite forms. On cooling, this transforms into a coarse prior austenite grain martensite. The grain size decreases with increasing distance from the fusion boundary [22] as the peak temperature attained decreased in the austenite region, resulting in a fine prior austenite grain martensite. When the thermal cycle peak temperature attained was in the range between AC1 and AC3 , only partial transformation to austenite took place during heating. Consequently, the microstructure after cooling is a mixture of austenite transformed product and untransformed ferrite. Untransformed ferrite gets overtempered during the thermal cycle. Below AC1 , there is no visible change in the base metal microstructure. Microhardness was taken across the weld joint in the same sequence as the metallographic examination. As depicted in Fig. 3, a gradient in the hardness values due to gradient in the microstructure is observed; a minimum in hardness value that corresponds to the ICHAZ microstructure is vivid.
Fig. 3. Microhardness variation across weld joint due to graded microstructure.
V. Shankar et al. / Materials Science and Engineering A 528 (2011) 8428–8437
8431
Fig. 4. Effect of temperature on the cyclic stress response curves for base metal and weld joint.
3.2. Effect of temperature on cyclic stress response, fatigue life and resultant microstructure The effect of temperature on the cyclic stress response curves (CSR) under continuous cycling condition for base metal and weld joint (WJ) are illustrated in Fig. 4. As the temperature is increased from ambient temperature to 823 K, the overall stress response values decreases by ∼225 MPa. Further decrease of ∼60 MPa in the overall stress response values is observed as the temperature is increased from 823 K to 873 K. The final load drop in the cyclic stress response occurs due to the initiation and propagation of fatigue cracks. Peak tensile stress at the half-life (i.e. at half of the number of cycles to failure) was taken as saturation or half-life stress and the cycle number corresponding to a drop of 20% from the half-life stress was taken as fatigue life [23]. A comparison of fatigue life at various applied strain amplitudes for both base metal and weld joint is given in Fig. 5(a) and (b) at 823 K and 873 K. It is observed that while at 823 K there is not much difference in the life values between the base metal and the weld joint (Fig. 5(a)), that at 873 K, the difference in the fatigue life values is quite noticeable (Fig. 5(b)). As observed in Fig. 4, the cyclic stress response of both base metal and the weld metal is similar, i.e. continuous softening from first cycle onwards. This is attributed to the microstructural degradation occurring during cycling in both base metal and the weld joint. However in the case of a sample containing weld joint, the response of each microstructural component to the applied external strain is different depending upon its mechanical properties. Hence the amount of softening which is directly linked to the cumulated plastic strain is not uniform in each microstructural component of the weld joint. The response observed during cycling in the form of hysteresis loops are the contributions from each component. Hence lower CSR of the weld joint compared to the CSR of the base metal is due to the presence of heterogeneous microstructure of varying mechanical properties and presence of a soft zone in the HAZ. Cyclic softening right from first cycle onwards in both base material and weld joint is observed and which agrees well with the reported observations typical of normalized and tempered ferritic/martensitic steels [24–27]. As mentioned earlier, Grade 91 steel is used in the normalized plus tempered condition and has a tempered martensitic structure. It derives its high temperature strength from several factors, among which are the lath structure, the high density of dislocations, sub-boundaries decorated with M23 C6 carbides, and fine Nb–V carbonitride precipitates of
Fig. 5. Fatigue life comparison between base metal and weld joint at 823 K (a) and 873 K (b).
MX type in the intragranular region. This microstructure is quite stable under high temperature exposure without deformation [28]. However, a decrease in hardness values has been reported due to microstructural changes under plastic deformation [28]. The lath structure which should remain stable under plastic deformation at high temperature, evolves towards equiaxed subgrain structure and significant dislocation-free areas appear. Also, coarsening of the strengthening carbides starts to occur. During fatigue deformation, the high number of dislocations present in the martensitic lath structure undergoes a back and forth movement under cycling [29] and they try to get rearranged into a lower energy configuration such as cells and subgrains causing cyclic softening in the material. During the process, many dislocations get annihilated and coarsening of carbides occur; this process is highly sensitive to temperature. Added to this, strain localization due to heterogeneity in microstructure and cavitation around coarsened carbides due to multiaxial stress occur in the weld joint. These together lead to differences in the cyclic stress response and fatigue lives of the BM and WJ. As depicted in Fig. 6(a) and (b), dislocation rearrangement occurs at ambient temperature as a result of cyclic deformation and dislocations move to the lath boundaries (Fig. 6(b)). This is considered to be the manifestation of interactions among mobile dislocations and the lath boundary dislocations resulting in cyclic softening. At high temperatures, thermally activated dislocation motion (climb/cross-slip) accelerates the annihilation mechanism and coarsens the laths and subgrains [24–27,30,31], and hence causes much higher cyclic softening. The resultant
8432
V. Shankar et al. / Materials Science and Engineering A 528 (2011) 8428–8437
Fig. 6. Effect of temperature on the resultant microstructure corresponding to the samples cycled at ambient ((a) and (b)) and 873 K (c) at ±0.6% strain amplitude up to failure.
microstructure due to fatigue deformation at high temperatures such as 873 K consists of larger subgrains (Fig. 6(c)) indicative of larger recovery. 3.3. Effect of application of hold on cyclic stress response and fatigue life The effect of application of strain hold at both 823 K and 873 K is an overall lowering of the peak tensile stress response compared to the continuous cycling cyclic stress response. A representative cyclic stress response curve of Grade 91 weld joint tested at 873 K and total strain amplitude of ±0.6% is depicted in Fig. 7(a). The overall peak tensile stress response under both tension and compression hold are much lower compared to the continuous cycling cyclic stress response. During the extended period of hold application, creep deformation occurs that results in stress relaxation. The elastic strain is converted partially into inelastic strain. A depiction of hysteresis loops for the first and half-life cycles for the 10 min TH test at 873 K is made in Fig. 7(b). The increased plastic strain under hold has been reported to have a correlation with cells and
sub-grain size [20]. The direction of hold has also been found to have a pronounced effect on the resultant microstructure such as the cell size. A comparatively larger and better defined cells under compression hold have been observed earlier in the base metal [20]. Based on these earlier observations on the base material [20], it may be postulated that creep in the form of larger recovery effects and also rapid coarsening of carbides result in a comparatively lower cyclic stress response of the hold tests performed on the Grade 91 weld joints. The effect of application of strain hold for durations ranging from 1 min to 30 min under peak tension and 1 min to 10 min under compression hold data on the fatigue life is compared with the continuous cycling fatigue life in Fig. 7(c). It vividly shows that the fatigue life reduce upon the application of either tensile or compressive hold. As the tensile hold duration is increased, the fatigue life decreases. Same observation was made earlier on the nearly homogeneous microstructure of the base metal and the reason for decease in fatigue life with increase in tensile hold duration in the base metal was ascribed to the increased amount of plastic strain accumulated during the extended period of hold application, larger
V. Shankar et al. / Materials Science and Engineering A 528 (2011) 8428–8437
8433
Fig. 8. Comparison of hardness profiles taken across the fusion line of the weld joints tested under continuous cycling and under the application of various durations of tensile (a) and compressive holds.
that microstructural changes such as substructural recovery and coarsening of strengthening carbides give rise to a decrease in the hardness value. Keeping this in mind, hardness profiles were taken across the fusion line for all the tested specimens. From Fig. 8(a) and (b)) it is clear that the hardness values decreases with an increase in temperature of testing and further with an application of hold in either of the directions. Also, as the duration of tensile hold was increased, the hardness value decreased, confirming that the hardness decrease is due to substructural changes and carbide coarsening. Detailed quantitative microstructural analysis of each zone of the weld joint in different test condition is being carried out to further quantify the above observations. Fig. 7. Effect of hold on the cyclic stress response (CSR) curves for base metal and weld joint at 873 K (a), representation of hysteresis loop generated during a 10 min tension hold test (b) and comparison of fatigue life under application of hold (c).
substructural changes and coarsening of carbides [20]. Application of a short duration 1 min compressive hold shows a much larger decrease in fatigue life compared to even longer duration tension hold up to 30 min. Compression hold being more damaging than tensile hold due to the deleterious effects of surface oxides in the base metal of this alloy is now well established [20,32–35] and the lower life of the weld joint is due to the same reason. This will be described in further detail in the next section. It is known
3.4. Failure location The fatigue tested samples were replicated to observe the origin of main crack that might have led to the final failure. Representative replicas along with the failed samples are shown in Fig. 9. It is observed that the failure occurs in the ICHAZ at 873 K under both continuous cycling and under compression hold whereas under tension hold, cracking is very close to the fusion line and is contained in the WM. This is observed consistently irrespective of the duration of hold in a given direction. In order to reconfirm the region/origin of failure, the longitudinal sections of all the failed samples tested under varied test conditions were analyzed. The
8434
V. Shankar et al. / Materials Science and Engineering A 528 (2011) 8428–8437
Fig. 9. Failure location for tests carried out at 873 K, ±0.6% under continuous cycling (a), compression hold (b) and tension hold (c). Respective replicas taken from the surface of failed samples are also shown adjacent to the failed samples.
Fig. 10. Microstructure obtained near the fractured surface for longitudinal section of specimens tested at room temperature (a) and at 873 K under continuous cycling (b), 10 min-CH (c) and 10 min-TH (d).
Fig. 11. Fractograph of specimen tested at 873 K under continuous cycling ((a) and (b)) and compression hold for 1 min ((c) and (d)); contour in the outer extremity of the fractured surface (marked by white dotted lines) marks localized plastic deformation.
V. Shankar et al. / Materials Science and Engineering A 528 (2011) 8428–8437
8435
Fig. 12. Specimen tested at 873 K under compression hold for 1 min shows bright field image as a result of dislocation rearrangements (a) and SEM images of the intercritical HAZ illustrating cavities around the coarsened M23 C6 carbides due to localized plastic deformation and triaxial state of stress (b) and (c).
microstructures obtained near the fractured surface for some of the representative samples are shown in Fig. 10(a)–(d). The microstructural analysis shows that whereas the failure mostly occurs in the intercritical region of the HAZ, some failures occur in the base material or within the weld metal. For example, whereas the sample failed in the parent base metal under ambient temperaturecontinuous cycling (Fig. 10(a)), that at 873 K, the failure location is in the ICHAZ under continuous cycling (Fig. 10(b)) and under compression hold (1 min hold and 10 min hold, Fig. 10(c)). However under tension hold (1, 10 and 30 min hold), the microstructure near the fractured surface is that of the weld metal (Fig. 10(d)). The fractured surface of the samples tested at 873 K under continuous cycling (Fig. 11(a) and (b)) and under 1 min compression
hold (Fig. 11(c) and (d)) are depicted, respectively. The surface topography indicates dark contour (marked by white dotted lines) on the periphery of the fractured surface (∼0.5 mm below the sample surface). Magnified image of the contour shows severe localized plastic deformation (Fig. 11(b)) under continuous cycling. Also, numerous creep cavities, separate or coalesced in these localized regions and their linkage leading to cracking in this dark contour region is observed under compression hold (Fig. 11(c) and (d)). This indicates that the contour could be the soft intercritical region in the HAZ. It may be stated that during each cycle, the plastic strain accumulated in different regions of the sample is different depending upon the mechanical properties of individual microstructures. During high temperature experiment, the soft zone of HAZ
Fig. 13. Large number of short secondary cracks perpendicular to loading direction under CH as compared to fewer unidirectional longer cracks under TH (a) is due to the role played by surface oxides during compression hold experiment (b).
8436
V. Shankar et al. / Materials Science and Engineering A 528 (2011) 8428–8437
Table 3 Chemical composition obtained by using EDAX taken across the contour (in wt%). Element
Base
At the dark contour
Weld
Al Si Nb Mo V Cr
1.25 1.04 0.29 0.66 0.42 8.96
1.25 1.04 0.29 0.66 0.42 8.96
0.19 0.6 0.66 1.45 0.57 8.55
accumulates larger plastic strain which is constrained by the surrounding material such as the comparatively stronger weld and base metal. Also much coarser and fewer M23 C6 carbides were observed to form in the ICHAZ (Fig. 12(a)–(c)) which is in line with that reported under pure creep conditions [7–10]. Goyal et al. [36] and Nonaka et al. [37] have shown the importance of multiaxiality of stress for creep cavitation around coarse precipitates in Cr–Mo steel weld joints. On similar understanding, under creep–fatigue interaction condition, the vital role played by multiaxiality of stress for cavity nucleation around the coarsened carbides in soft ICHAZ cannot be ruled out. Vacancies generated during fatigue enhance creep climb and carbide coarsening that assist in cavity nucleation and hence enhanced creep deformation occur during creep–fatigue interaction conditions. EDAX analysis was carried out across the contour observed in the fractographs. The chemical composition showed that the contour was a part of the base metal and not the weld metal. The chemical composition is given in Table 3. Hence based on their location and chemical composition analysis, the regions A, B and C marked in Fig. 11(d) were identified as base metal, HAZ and WM, respectively. The more damaging effect of application of compression hold on life reduction compared to the tension hold may be explained on the basis of the observations made on the polished surface of the longitudinal section of the failed samples tested under tension hold and compression hold at 873 K (Fig. 13). These micrographs show multiple crack initiation sites and oxidation assisted crack initiation and propagation for sample tested under compression hold. Also, the number of secondary cracks found under compression hold is much higher than that found under tension hold. The larger compression dwell sensitivity, i.e. compression hold being more damaging than tension hold in Grade 91 steel essentially arises from the deleterious effect of oxidation [20,32–35]. Oxidation assisted crack initiation and propagation is vividly shown in Fig. 13(b). Brittle surface oxides forms on the sample surface during hold at high temperatures break upon unloading and further cycling in tensile direction after removal of compression hold. Oxygen impregnates through the freshly exposed metal layer causing more damage in the material. This explains the compression dwell sensitivity of the alloy as previously reported [20,32–35]. Hence the decrease in fatigue life of a weld joint by strain hold arises from the acceleration of crack initiation due to creep strain concentration in the HAZ as well as the acceleration of crack growth due to creep voids caused by plastic restraint in HAZ. The larger compressive dwell sensitivity is largely due to the deleterious effect of oxidation. 4. Conclusions Both continuous cycling and hold time tests on Grade 91 steel exhibits a continuous cyclic softening in both the base metal and the weld joint. The high dislocation density initially present in the martensitic lath structure undergoes a rearrangement into a lower energy configuration such as cells and subgrains causing cyclic softening in the material and this process is highly sensitive to temperature and application of hold. The lower cyclic stress response curve of the weld joint compared to the base metal is attributed to
the presence of a heterogeneous microstructure and the soft intercritical heat affected zone of the weld joint. Failure location in the weld joint is found to depend upon temperature and application of hold. While at ambient temperature, failure occurs in the base metal, at high temperatures, failure arises in the ICHAZ under continuous cycling and compression hold. The failure location is very close to the fusion line and is contained in the WM under tensile hold. Strain localization in the soft zone of the heat affected zone, sub-surface creep cavity formation in this region and their linkage causes enhanced crack propagation that translates into lower fatigue life of the weld joint at high temperatures. The compression dwell sensitivity is largely due to the deleterious effect of the surface oxides. Acknowledgements The authors are grateful to Dr Baldev Raj, former Director, IGCAR, Sri S.C. Chetal, Director, IGCAR and Dr T. Jayakumar, Group Director Metallurgy and Materials Group, IGCAR, Kalpakkam for their support and encouragement during this research. The authors also thank Mrs. M. Radhika, PMD, IGCAR for carrying out SEM investigation and Sri P. Sukumar, NDED for taking the replicas. The experimental assistance during fatigue testing by Sri K. Marriappan during this study is gratefully acknowledged. References [1] J.M. Vitek, R.L. Klueh, Metall. Mater. Trans. A 14 (1983) 1047–1055. [2] B.W. Jones, C.R. Hills, D.H. Polonis, Metall. Trans. A 22 (1991) 1049–1058. [3] I.A. Shibli, in: A. Strang, R.D. Conroy, W.M. Banks, M. Blackler, J. Leggett, G.M. McColvin, S. Simpson, M. Smith, F. Star, R.W. Vanstone (Eds.), Proceedings of the Sixth International Charles Parsons Turbine Conference, 16–18 September 2003, Trinity College Dublin, Ireland, 2003, pp. 261–279. [4] C. Middleton, E. Metcalfe, IMechE Proceedings, London, UK, 1990 (Paper C386/027). [5] Shibli, Proceedings of the Swansea Creep Conference, University of Swansea and EPRI, Swansea, UK, 2001. [6] J.A. Francis, W. Mazur, H.K.D.H. Bhadeshia, Mater. Sci. Technol. 22 (12) (2006) 1387–1395. [7] H. Cerjak, E. Letofsky, in: R. Viswanathan, J. Nutting (Eds.), Advanced Heat Resistance Steels for Power Generation, 1998, pp. 611–621. [8] J.M. Brear, A. Fairman, C.J. Middleton, L. Polding, Key Eng. Mater. 171–174 (3) (2000) 5–42. [9] C.J. Middleton, J.M. Brear, R. Munson, R. Vishwanathan, in: R. Vishwanathan, W.T. Bakker, J.D. Parker (Eds.), Proceedings of the 3rd Conference on Advances in Materials Technology for Fossil Power Plant, The Inst. Mater., London, 2001, pp. 69–78. [10] E. Letofsky, H. Cerjak, I. Papst, P. Warbichler, in: R. Vishwanathan, W.T. Bakker, J.D. Parker (Eds.), Proceedings of the 3rd Conference on Advances in Materials Technology for Fossil Power Plant, The Inst. Mater., London, 2001, pp. 133–142. [11] F. Masuyama, M. Matsui, N. Komai, Key Eng. Mater. 171–174 (2000) 99–108. [12] M. Matsui, M. Tabuchi, T. Watanabe, K. Kubo, J. Kinugawa, F. Abe, ISIJ Int. 41 (2001) S126–S130. [13] M. Tabuchi, T. Watanabe, K. Kubo, M. Matsui, J. Kinugawa, F. Abe, Int. J. Press. Vessels Piping 78 (2001) 779–784. [14] M. Tabuchi, M. Matsui, T. Watanabe, H. Hongo, K. Kubo, F. Abe, Mater. Sci. Res. Int. 9 (2003) 23–28. [15] F. Vivier, Fluage a’ 500 ◦ C d’ un joint sounde’ d’ un acier 9Cr–1Mo modifie’. Evolution de la microstructure & Comportement me’canique, PhD thesis, 2009, Ecole des Mines de Paris, France. [16] D. Jandova, J. Kasl, V. Kanta, Proceedings of the 2nd International ECCC Conference, Empa, CH-8600 Dübendorf Akademie, Überlandstrasse 129, April, 2009. [17] V. Gaffard, Experimental study and modeling of high temperature creep flow and damage behaviour of 9Cr–1Mo–NbV steel weldments, PhD thesis, 2004, Ecole des Mines de Paris, France. [18] T. Watanbe, M. Tabuchi, M. Yamazaki, H. Hongo, T. Tanabe, Int. J. Press. Vessels Piping 83 (2006) 63–71. [19] ASTM E606-92, Standard recommended practice for constant-amplitude lowcycle fatigue testing, Annual Book of ASTM Standards, 1994, 03.01, pp. 522–536. [20] V. Shankar, M. Valsan, K. Bhanu Sankara Rao, R. Kannan, S.L. Mannan, S.D. Pathak, Mater. Sci. Eng. A 437 (2) (2006) 413–422. [21] K.S. Chandravathi, K. Laha, K.B.S. Rao, S.L. Mannan, Mater. Sci. Technol. 17 (2001) 559–565. [22] D.P. Edmonds, B. Chew, Effect of chemical composition on weld microstructure in some commercial 9Cr–1Mo steel boiler tubes. Report TPRD/M/1258/N82. Central Electricity Generating Board, Marchwood, 1982. [23] K. Bhanu Sankara Rao, M. Valsan, R. Sandhya, S.K. Ray, S.L. Mannan, P. Rodriguez, Int. J. Fatigue 7 (3) (1985) 141–147.
V. Shankar et al. / Materials Science and Engineering A 528 (2011) 8428–8437 [24] S. Kim, J.R. Weertman, Metall. Trans. A 19 (1988) 999–1007. [25] B. Fournier, M. Sauzay, A. Renault, F. Barcelo, A. Pineau, J. Nucl. Mater. 386–388 (386) (2009) 71–80. [26] H. Li, A. Nishimura, T. Nagasaka, T. Muroga, Fusion Eng. Des. 81 (2006) 2907–2912. [27] A.F. Armas, C. Petersen, R. Schmitt, M. Avalos, I. Alvarez-Armas, J. Nucl. Mater. 307–311 (307) (2002) 509–513. [28] S.T. Kimmins, M.C. Coleman, D.J. Smith, in: B. Wilshire, R.W. Evans (Eds.), Proceedings of the Fifth International Conference on Creep and Fatigue of Engineering Materials and Structure, Swansea, 1993. [29] D.C. Lord, L.F. Coffin Jr., Metall. Trans. 4 (1973) 1647–1654. [30] W.B. Jones, in: A.K. Khare (Ed.), Proceedings of the ASM International Conference on Production, Fabrication, Properties and Application of Ferritic Steels for High Temperature Applications, Warren, PA, 1981, pp. 221–235.
8437
[31] P. Marmy, T. Kruml, J. Nucl. Mater. 377 (2008) 52–58. [32] B. Fournier, M. Sauzay, C. Caës, M. Noblecourt, M. Mottot, A. Bougault, V. Rabeau, J. Man, O. Gillia, P. Lemoine, A. Pineau, Int. J. Fatigue 30 (October–November (10–11)) (2008) 1797–1812. [33] J.S. Dubey, H. Chilukuru, J.K. Chakravartty, M. Schwienheer, A. Scholz, W. Blum, Mater. Sci. Eng. A 406 (October (1–2)) (2005) 152–159. [34] B. Fournier, M. Sauzay, C. Caës, M. Noblecourt, M. Mottot, A. Bougault, V. Rabeau, A. Pineau, Int. J. Fatigue 30 (April (4)) (2008) 649–662. [35] B. Fournier, M. Sauzay, C. Caës, M. Noblecourt, M. Mottot, A. Bougault, V. Rabeau, A. Pineau, Int. J. Fatigue 30 (April (4)) (2008) 663–676. [36] S. Goyal, K. Laha, K.S. Chandravathi, P. Parameswaran, M.D. Mathew, Philos. Magazine 91 (23) (2011) 3128–3154. [37] I. Nonaka, S. Torihata, S. Kihara, H. Umaki, Mater. High Temp. 15 (2) (1998) 69–73.
Contents lists available at ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Creep–fatigue-oxidation interaction in Grade 91 steel weld joints for high temperature applications Vani Shankar ∗ , R. Sandhya, M.D. Mathew Mechanical Metallurgy Division, Indira Gandhi Centre for Atomic Research, Kalpakkam 603102, Tamil Nadu, India
a r t i c l e
i n f o
Article history: Received 3 February 2011 Received in revised form 12 July 2011 Accepted 22 July 2011 Available online 29 July 2011 Keywords: Grade 91 steel weld joint Type IV cracking Low cycle fatigue Creep–fatigue interaction Oxidation
a b s t r a c t The weld joint consisting of a heterogeneous microstructure exhibits a lower fatigue life than that of the base metal. This is due to the presence of soft zone in the heat affected zone (HAZ). At high temperatures and under hold application, strain localization occurs in the soft intercritical HAZ (ICHAZ). Sub-surface creep cavity formation in the soft region and their linkage causes enhanced crack propagation and that translates into lower fatigue life of the weld joint at high temperatures. Occurrence of compression dwell sensitivity in the material is attributed to the presence of surface oxides. © 2011 Elsevier B.V. All rights reserved.
1. Introduction Modified version of plain 9Cr–1Mo steel in which Nb and V are optimized (Grade 91 steel) are being extensively used for super heater tubing, headers and piping of conventional as well as nuclear power plants with steam temperatures up to 866 K [1,2]. Compared to the austenitic stainless steels, the alloy has better resistance to thermal fatigue and stress corrosion cracking. The Grade 91 steel has better monotonic tensile, creep and fatigue properties at elevated temperatures compared to plain 9Cr–1Mo steel. Steam generators (SG) are complex large structures that are made in parts and integrated by welding techniques. Welds are weak links in any structure and many failures in high temperature components are reported to be weld-related [3–5]. During welding, the temperature sensitive microstructure of Grade 91 steel gets modified to various extents depending upon the temperature it is exposed to during the weld thermal cycles; highest temperature is reached by HAZ closest to the weld metal. Hence the narrow band of heat affected zone (HAZ) is a complex variation of microstructure with varied mechanical properties. Many premature failures at weld joints (WJ) have been identified to occur in a narrow fine grained or inter critical region of the HAZ in stronger 9–12 Cr steels such as Grade 91 ferritic steel [3–14] and they are termed as type IV failures.
∗ Corresponding author. Tel.: +91 44 274800x21147; fax: +91 44 27480075. E-mail address: [email protected] (V. Shankar). 0921-5093/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2011.07.046
In sodium cooled fast reactors (SFRs), the components operating at high temperatures are subjected to heating and cooling transients (sharp thermal gradients from surface to core of thick components) during startup and shutdown. This onsets thermal stresses in material if thermal strain (expansion/contraction) is totally/partly constrained. Repetition of these thermal transient results in strain-controlled low cycle fatigue of the components. In addition, steady state operation at elevated temperatures introduces creep, resulting in creep–fatigue interaction condition. In order to have more reliability in the performance of a component, it is essential to understand the deformation mechanisms operative under pure fatigue, pure creep and creep–fatigue interaction conditions. Creep experiments on weld joints of Grade 91 steel have shown that fracture usually occurs in the base metal in the short term tests at high stresses, whereas under long term creep conditions at low stresses Type IV failure mode appears [7–10]. In high chromium ferritic martensitic steels, significant reduction in rupture life as compared to the specimens without weld joint have been reported due to Type IV failure [6,9–14] especially at low stresses. Under creep at high stresses, failure has been reported to occur in the weld metal [15–18]. Whereas the creep performance of weld joints of Grade 91 steel is much better understood, the low cycle fatigue (LCF) and creep–fatigue interaction (CFI) behavior of the weld joints are not so well established. Hence the aim of the present paper is to study creep–fatigue interaction (CFI) behavior, identify the underlying deformation mechanism and correlate with the failure mode occurring in Grade 91 steel weld joints.
V. Shankar et al. / Materials Science and Engineering A 528 (2011) 8428–8437
8429
Table 1 Chemical composition (wt%) of base metal and weld metal. Element
C
Cr
Mo
Ni
V
Nb
N
S
P
Cu
Co
Base metal Weld metal
0.11 0.1
9.3 9
0.99 1.0
0.14 0.7
0.25 0.17
0.1 0.06
0.068 0.055
0.008 0.012
0.020 0.009
<0.05
0.016
2. Experimental Chemical composition of the 30 mm thick rolled plate used for current study and the weld metal is given in Table 1. Due to the high reliability required for nuclear components, the material has closely controlled composition, low residual element concentration and low inclusion content. The base material was given a normalizing treatment at 1313 K for 1 h followed by air cooling and subsequently tempering at 1033 K for 1 h followed by air cooling, before machining the fatigue specimens. Microstructural characterizations were carried out using optical microscope, scanning and transmission electron microscopes (TEM). Samples for the optical metallography were etched using Vilella’s reagent (1 g of picric acid + 5 ml conc. HCl + 100 ml ethyl alcohol). Samples for TEM were prepared by mechanically thinning the foils to 100 m thickness and finally electrolytic thinning at 20 V using stainless steel electrodes in a double jet electropolisher. The electrolyte used was 20% perchloric acid in methanol at 233 K. For fabricating LCF specimens containing weld joint, the plates were joined along the rolling direction by shielded metal arc welding (SMAW) process using voltage and current approximately 20 V and 100 A, respectively. A double-V configuration, with an included angle of 70◦ , a root face of 2 mm, and a root gap of 3.15 mm was used. Multiple passes were employed to fill the groove. An inter-pass temperature of 423 K was maintained during welding. Matching filler wire was used for welding. The weld pads were examined by radiography for their soundness. A nearly symmetrical weld profile was obtained after multipass welding (Fig. 1(a)). Specimens were taken from the centre of the welded plate (Fig. 1(b)). A schematic of specimen geometry containing the weld joint is shown in Fig. 1(c). The weld joint consists of three regions—base metal (BM), heat affected zone (HAZ) and weld metal (WM). The standard LCF specimen was selected in such a way that the WM was in the centre of the gauge length and the two HAZs were in the two sides of the WM and the rest of the sample was BM. Using metallographic and non destructive techniques, the double V weld profile, different zones of the HAZ and the amount of BM contained in the sample was first established. The maximum width of WM was around 10 mm with ∼3.5–4 mm HAZ and 3 mm BM on each side of WM. Bars of 110 mm length and 25 mm × 25 mm square cross section were cut from the welded plate. Post welding heat treatment (PWHT) was given to the bars containing weld joints at 1033 K for 3 h followed by air cooling. A nearly symmetrical weld profile was obtained (Fig. 1(b)). Low cycle fatigue (LCF) and creep–fatigue interaction (CFI) tests were conducted in air, under fully reversed, total axial strain control mode in accordance with ASTM specification E606 [19] in a
closed loop servo hydraulic testing system equipped with a radiant heating furnace. The temperature variation along the gauge length of the specimen did not exceed ±2 K. Continuous cycling tests were performed using a triangular waveform and hold time experiments were carried out using a trapezoidal waveform. LCF tests under continuous cycling condition were carried out at room temperature at strain amplitudes of ±0.6% and at 873 K at strain amplitude ±0.25%, ±0.4%, ±0.6% and ±1%. CFI tests were conducted by introducing hold of 1, 10 and 30 min at peak tension and 1 and 10 min in peak compression under ±0.6% strain amplitude. All the tests were carried out at a constant strain rate of 3 × 10−3 s−1 . A table summarizing the testing conditions is given as Table 2. After fatigue tests, the failed samples were pulled out to separate the fractured surface. Fractographic analysis was performed using SEM after which the longitudinal sections of the two separated parts were analyzed using an optical microscope equipped with an image analyzer. Microstructure near the fractured surface was studied for all the tested specimens. Microhardness was taken across the fusion line using 200 gmf load in a microhardness tester. Since failure usually occurred in one of the two sides of the weld joint (WJ), the microhardness measurements was performed in the side that had not cracked.
3. Results and discussion 3.1. Initial characterization of material The base material after normalizing and tempering treatment reveal a tempered martensitic microstructure with extensive precipitation of M23 C6 at the boundaries and MX type precipitates within the matrix [20]. Starting from weld metal towards base material, the microstructure is as follows: weld metal (Fig. 2(a)) with a tempered martensitic structure, tempered martensite structure with islands of ␦-ferrite (Fig. 2(b)), fine grained region (Fig. 2(c)), inter-critical structure (Fig. 2(d)) and unaffected base material (Fig. 2(e)) with a tempered martensitic structure. The total width of the HAZ is ∼3–4 mm. The intercritical region is approximately 2.5–3 mm away from the fusion line and close to BM. The region closest to the fusion boundary experiences a peak temperature above AC4 (the boundary between ␥ and ␥ + ␦ phase fields). At this high temperature, ␦ ferrite form along the austenite grain boundaries. The high alloy content of the steel resulting from dissolution of carbonitrides [21], coupled with the high cooling rate in the weld thermal cycle causes the retention of some ␦ ferrite on cooling. The resulting structure after cooling is therefore
Fig. 1. Weld profile obtained after multipass welding (a), details of weld joint specimens taken from welded plate (b) and schematic of specimen geometry containing the weld joint (c).
8430
V. Shankar et al. / Materials Science and Engineering A 528 (2011) 8428–8437
Table 2 Summary of results showing effect of temperature, strain amplitude and application of hold on LCF properties. Temperature (K)
Strain amplitude (%)
Type and duration of hold
1st cycle stress (MPa)
Peak stress at half-life (MPa)
300
0.6
0
572
520
1030
0 0 0
823
0.25 0.4 0.6 (BM) 0.6 1.0 0.6 0.6
328 375 456 392 410 354 389
262 291 315 300 296 271 307
3206 1007 580 475 209 374 240
0.25 0.4 0.6 (BM) 0.6 1.0 0.6 0.6 0.6 0.6
0 0 0
299 340 386 360 322 331 345 340 342
220 239 260 240 245 212 237 251 232
2250 1218 500 534 223 510 400 235 324
873
0 1 min TH 1 min CH
0 1 min TH 10 min TH 1 min CH 10 min CH
Fatigue life
Fig. 2. Resultant graded microstructure in the Grade 91 steel weld joint due to thermal exposure to different temperatures during weld thermal cycles; optical micrographs of different regions of weld joint such as weld metal (a), coarse grain with ␦ ferrite (b), fine grained region (c), intercritical region (d) and base metal (e).
coarse prior austenite grain martensite with ␦ ferrite along the prior austenite grain boundaries. Further away from the fusion boundary, where the peak temperature was well above AC3 but below AC4 , the carbides that impede austenite grain growth get dissolved and hence coarse grain austenite forms. On cooling, this transforms into a coarse prior austenite grain martensite. The grain size decreases with increasing distance from the fusion boundary [22] as the peak temperature attained decreased in the austenite region, resulting in a fine prior austenite grain martensite. When the thermal cycle peak temperature attained was in the range between AC1 and AC3 , only partial transformation to austenite took place during heating. Consequently, the microstructure after cooling is a mixture of austenite transformed product and untransformed ferrite. Untransformed ferrite gets overtempered during the thermal cycle. Below AC1 , there is no visible change in the base metal microstructure. Microhardness was taken across the weld joint in the same sequence as the metallographic examination. As depicted in Fig. 3, a gradient in the hardness values due to gradient in the microstructure is observed; a minimum in hardness value that corresponds to the ICHAZ microstructure is vivid.
Fig. 3. Microhardness variation across weld joint due to graded microstructure.
V. Shankar et al. / Materials Science and Engineering A 528 (2011) 8428–8437
8431
Fig. 4. Effect of temperature on the cyclic stress response curves for base metal and weld joint.
3.2. Effect of temperature on cyclic stress response, fatigue life and resultant microstructure The effect of temperature on the cyclic stress response curves (CSR) under continuous cycling condition for base metal and weld joint (WJ) are illustrated in Fig. 4. As the temperature is increased from ambient temperature to 823 K, the overall stress response values decreases by ∼225 MPa. Further decrease of ∼60 MPa in the overall stress response values is observed as the temperature is increased from 823 K to 873 K. The final load drop in the cyclic stress response occurs due to the initiation and propagation of fatigue cracks. Peak tensile stress at the half-life (i.e. at half of the number of cycles to failure) was taken as saturation or half-life stress and the cycle number corresponding to a drop of 20% from the half-life stress was taken as fatigue life [23]. A comparison of fatigue life at various applied strain amplitudes for both base metal and weld joint is given in Fig. 5(a) and (b) at 823 K and 873 K. It is observed that while at 823 K there is not much difference in the life values between the base metal and the weld joint (Fig. 5(a)), that at 873 K, the difference in the fatigue life values is quite noticeable (Fig. 5(b)). As observed in Fig. 4, the cyclic stress response of both base metal and the weld metal is similar, i.e. continuous softening from first cycle onwards. This is attributed to the microstructural degradation occurring during cycling in both base metal and the weld joint. However in the case of a sample containing weld joint, the response of each microstructural component to the applied external strain is different depending upon its mechanical properties. Hence the amount of softening which is directly linked to the cumulated plastic strain is not uniform in each microstructural component of the weld joint. The response observed during cycling in the form of hysteresis loops are the contributions from each component. Hence lower CSR of the weld joint compared to the CSR of the base metal is due to the presence of heterogeneous microstructure of varying mechanical properties and presence of a soft zone in the HAZ. Cyclic softening right from first cycle onwards in both base material and weld joint is observed and which agrees well with the reported observations typical of normalized and tempered ferritic/martensitic steels [24–27]. As mentioned earlier, Grade 91 steel is used in the normalized plus tempered condition and has a tempered martensitic structure. It derives its high temperature strength from several factors, among which are the lath structure, the high density of dislocations, sub-boundaries decorated with M23 C6 carbides, and fine Nb–V carbonitride precipitates of
Fig. 5. Fatigue life comparison between base metal and weld joint at 823 K (a) and 873 K (b).
MX type in the intragranular region. This microstructure is quite stable under high temperature exposure without deformation [28]. However, a decrease in hardness values has been reported due to microstructural changes under plastic deformation [28]. The lath structure which should remain stable under plastic deformation at high temperature, evolves towards equiaxed subgrain structure and significant dislocation-free areas appear. Also, coarsening of the strengthening carbides starts to occur. During fatigue deformation, the high number of dislocations present in the martensitic lath structure undergoes a back and forth movement under cycling [29] and they try to get rearranged into a lower energy configuration such as cells and subgrains causing cyclic softening in the material. During the process, many dislocations get annihilated and coarsening of carbides occur; this process is highly sensitive to temperature. Added to this, strain localization due to heterogeneity in microstructure and cavitation around coarsened carbides due to multiaxial stress occur in the weld joint. These together lead to differences in the cyclic stress response and fatigue lives of the BM and WJ. As depicted in Fig. 6(a) and (b), dislocation rearrangement occurs at ambient temperature as a result of cyclic deformation and dislocations move to the lath boundaries (Fig. 6(b)). This is considered to be the manifestation of interactions among mobile dislocations and the lath boundary dislocations resulting in cyclic softening. At high temperatures, thermally activated dislocation motion (climb/cross-slip) accelerates the annihilation mechanism and coarsens the laths and subgrains [24–27,30,31], and hence causes much higher cyclic softening. The resultant
8432
V. Shankar et al. / Materials Science and Engineering A 528 (2011) 8428–8437
Fig. 6. Effect of temperature on the resultant microstructure corresponding to the samples cycled at ambient ((a) and (b)) and 873 K (c) at ±0.6% strain amplitude up to failure.
microstructure due to fatigue deformation at high temperatures such as 873 K consists of larger subgrains (Fig. 6(c)) indicative of larger recovery. 3.3. Effect of application of hold on cyclic stress response and fatigue life The effect of application of strain hold at both 823 K and 873 K is an overall lowering of the peak tensile stress response compared to the continuous cycling cyclic stress response. A representative cyclic stress response curve of Grade 91 weld joint tested at 873 K and total strain amplitude of ±0.6% is depicted in Fig. 7(a). The overall peak tensile stress response under both tension and compression hold are much lower compared to the continuous cycling cyclic stress response. During the extended period of hold application, creep deformation occurs that results in stress relaxation. The elastic strain is converted partially into inelastic strain. A depiction of hysteresis loops for the first and half-life cycles for the 10 min TH test at 873 K is made in Fig. 7(b). The increased plastic strain under hold has been reported to have a correlation with cells and
sub-grain size [20]. The direction of hold has also been found to have a pronounced effect on the resultant microstructure such as the cell size. A comparatively larger and better defined cells under compression hold have been observed earlier in the base metal [20]. Based on these earlier observations on the base material [20], it may be postulated that creep in the form of larger recovery effects and also rapid coarsening of carbides result in a comparatively lower cyclic stress response of the hold tests performed on the Grade 91 weld joints. The effect of application of strain hold for durations ranging from 1 min to 30 min under peak tension and 1 min to 10 min under compression hold data on the fatigue life is compared with the continuous cycling fatigue life in Fig. 7(c). It vividly shows that the fatigue life reduce upon the application of either tensile or compressive hold. As the tensile hold duration is increased, the fatigue life decreases. Same observation was made earlier on the nearly homogeneous microstructure of the base metal and the reason for decease in fatigue life with increase in tensile hold duration in the base metal was ascribed to the increased amount of plastic strain accumulated during the extended period of hold application, larger
V. Shankar et al. / Materials Science and Engineering A 528 (2011) 8428–8437
8433
Fig. 8. Comparison of hardness profiles taken across the fusion line of the weld joints tested under continuous cycling and under the application of various durations of tensile (a) and compressive holds.
that microstructural changes such as substructural recovery and coarsening of strengthening carbides give rise to a decrease in the hardness value. Keeping this in mind, hardness profiles were taken across the fusion line for all the tested specimens. From Fig. 8(a) and (b)) it is clear that the hardness values decreases with an increase in temperature of testing and further with an application of hold in either of the directions. Also, as the duration of tensile hold was increased, the hardness value decreased, confirming that the hardness decrease is due to substructural changes and carbide coarsening. Detailed quantitative microstructural analysis of each zone of the weld joint in different test condition is being carried out to further quantify the above observations. Fig. 7. Effect of hold on the cyclic stress response (CSR) curves for base metal and weld joint at 873 K (a), representation of hysteresis loop generated during a 10 min tension hold test (b) and comparison of fatigue life under application of hold (c).
substructural changes and coarsening of carbides [20]. Application of a short duration 1 min compressive hold shows a much larger decrease in fatigue life compared to even longer duration tension hold up to 30 min. Compression hold being more damaging than tensile hold due to the deleterious effects of surface oxides in the base metal of this alloy is now well established [20,32–35] and the lower life of the weld joint is due to the same reason. This will be described in further detail in the next section. It is known
3.4. Failure location The fatigue tested samples were replicated to observe the origin of main crack that might have led to the final failure. Representative replicas along with the failed samples are shown in Fig. 9. It is observed that the failure occurs in the ICHAZ at 873 K under both continuous cycling and under compression hold whereas under tension hold, cracking is very close to the fusion line and is contained in the WM. This is observed consistently irrespective of the duration of hold in a given direction. In order to reconfirm the region/origin of failure, the longitudinal sections of all the failed samples tested under varied test conditions were analyzed. The
8434
V. Shankar et al. / Materials Science and Engineering A 528 (2011) 8428–8437
Fig. 9. Failure location for tests carried out at 873 K, ±0.6% under continuous cycling (a), compression hold (b) and tension hold (c). Respective replicas taken from the surface of failed samples are also shown adjacent to the failed samples.
Fig. 10. Microstructure obtained near the fractured surface for longitudinal section of specimens tested at room temperature (a) and at 873 K under continuous cycling (b), 10 min-CH (c) and 10 min-TH (d).
Fig. 11. Fractograph of specimen tested at 873 K under continuous cycling ((a) and (b)) and compression hold for 1 min ((c) and (d)); contour in the outer extremity of the fractured surface (marked by white dotted lines) marks localized plastic deformation.
V. Shankar et al. / Materials Science and Engineering A 528 (2011) 8428–8437
8435
Fig. 12. Specimen tested at 873 K under compression hold for 1 min shows bright field image as a result of dislocation rearrangements (a) and SEM images of the intercritical HAZ illustrating cavities around the coarsened M23 C6 carbides due to localized plastic deformation and triaxial state of stress (b) and (c).
microstructures obtained near the fractured surface for some of the representative samples are shown in Fig. 10(a)–(d). The microstructural analysis shows that whereas the failure mostly occurs in the intercritical region of the HAZ, some failures occur in the base material or within the weld metal. For example, whereas the sample failed in the parent base metal under ambient temperaturecontinuous cycling (Fig. 10(a)), that at 873 K, the failure location is in the ICHAZ under continuous cycling (Fig. 10(b)) and under compression hold (1 min hold and 10 min hold, Fig. 10(c)). However under tension hold (1, 10 and 30 min hold), the microstructure near the fractured surface is that of the weld metal (Fig. 10(d)). The fractured surface of the samples tested at 873 K under continuous cycling (Fig. 11(a) and (b)) and under 1 min compression
hold (Fig. 11(c) and (d)) are depicted, respectively. The surface topography indicates dark contour (marked by white dotted lines) on the periphery of the fractured surface (∼0.5 mm below the sample surface). Magnified image of the contour shows severe localized plastic deformation (Fig. 11(b)) under continuous cycling. Also, numerous creep cavities, separate or coalesced in these localized regions and their linkage leading to cracking in this dark contour region is observed under compression hold (Fig. 11(c) and (d)). This indicates that the contour could be the soft intercritical region in the HAZ. It may be stated that during each cycle, the plastic strain accumulated in different regions of the sample is different depending upon the mechanical properties of individual microstructures. During high temperature experiment, the soft zone of HAZ
Fig. 13. Large number of short secondary cracks perpendicular to loading direction under CH as compared to fewer unidirectional longer cracks under TH (a) is due to the role played by surface oxides during compression hold experiment (b).
8436
V. Shankar et al. / Materials Science and Engineering A 528 (2011) 8428–8437
Table 3 Chemical composition obtained by using EDAX taken across the contour (in wt%). Element
Base
At the dark contour
Weld
Al Si Nb Mo V Cr
1.25 1.04 0.29 0.66 0.42 8.96
1.25 1.04 0.29 0.66 0.42 8.96
0.19 0.6 0.66 1.45 0.57 8.55
accumulates larger plastic strain which is constrained by the surrounding material such as the comparatively stronger weld and base metal. Also much coarser and fewer M23 C6 carbides were observed to form in the ICHAZ (Fig. 12(a)–(c)) which is in line with that reported under pure creep conditions [7–10]. Goyal et al. [36] and Nonaka et al. [37] have shown the importance of multiaxiality of stress for creep cavitation around coarse precipitates in Cr–Mo steel weld joints. On similar understanding, under creep–fatigue interaction condition, the vital role played by multiaxiality of stress for cavity nucleation around the coarsened carbides in soft ICHAZ cannot be ruled out. Vacancies generated during fatigue enhance creep climb and carbide coarsening that assist in cavity nucleation and hence enhanced creep deformation occur during creep–fatigue interaction conditions. EDAX analysis was carried out across the contour observed in the fractographs. The chemical composition showed that the contour was a part of the base metal and not the weld metal. The chemical composition is given in Table 3. Hence based on their location and chemical composition analysis, the regions A, B and C marked in Fig. 11(d) were identified as base metal, HAZ and WM, respectively. The more damaging effect of application of compression hold on life reduction compared to the tension hold may be explained on the basis of the observations made on the polished surface of the longitudinal section of the failed samples tested under tension hold and compression hold at 873 K (Fig. 13). These micrographs show multiple crack initiation sites and oxidation assisted crack initiation and propagation for sample tested under compression hold. Also, the number of secondary cracks found under compression hold is much higher than that found under tension hold. The larger compression dwell sensitivity, i.e. compression hold being more damaging than tension hold in Grade 91 steel essentially arises from the deleterious effect of oxidation [20,32–35]. Oxidation assisted crack initiation and propagation is vividly shown in Fig. 13(b). Brittle surface oxides forms on the sample surface during hold at high temperatures break upon unloading and further cycling in tensile direction after removal of compression hold. Oxygen impregnates through the freshly exposed metal layer causing more damage in the material. This explains the compression dwell sensitivity of the alloy as previously reported [20,32–35]. Hence the decrease in fatigue life of a weld joint by strain hold arises from the acceleration of crack initiation due to creep strain concentration in the HAZ as well as the acceleration of crack growth due to creep voids caused by plastic restraint in HAZ. The larger compressive dwell sensitivity is largely due to the deleterious effect of oxidation. 4. Conclusions Both continuous cycling and hold time tests on Grade 91 steel exhibits a continuous cyclic softening in both the base metal and the weld joint. The high dislocation density initially present in the martensitic lath structure undergoes a rearrangement into a lower energy configuration such as cells and subgrains causing cyclic softening in the material and this process is highly sensitive to temperature and application of hold. The lower cyclic stress response curve of the weld joint compared to the base metal is attributed to
the presence of a heterogeneous microstructure and the soft intercritical heat affected zone of the weld joint. Failure location in the weld joint is found to depend upon temperature and application of hold. While at ambient temperature, failure occurs in the base metal, at high temperatures, failure arises in the ICHAZ under continuous cycling and compression hold. The failure location is very close to the fusion line and is contained in the WM under tensile hold. Strain localization in the soft zone of the heat affected zone, sub-surface creep cavity formation in this region and their linkage causes enhanced crack propagation that translates into lower fatigue life of the weld joint at high temperatures. The compression dwell sensitivity is largely due to the deleterious effect of the surface oxides. Acknowledgements The authors are grateful to Dr Baldev Raj, former Director, IGCAR, Sri S.C. Chetal, Director, IGCAR and Dr T. Jayakumar, Group Director Metallurgy and Materials Group, IGCAR, Kalpakkam for their support and encouragement during this research. The authors also thank Mrs. M. Radhika, PMD, IGCAR for carrying out SEM investigation and Sri P. Sukumar, NDED for taking the replicas. The experimental assistance during fatigue testing by Sri K. Marriappan during this study is gratefully acknowledged. References [1] J.M. Vitek, R.L. Klueh, Metall. Mater. Trans. A 14 (1983) 1047–1055. [2] B.W. Jones, C.R. Hills, D.H. Polonis, Metall. Trans. A 22 (1991) 1049–1058. [3] I.A. Shibli, in: A. Strang, R.D. Conroy, W.M. Banks, M. Blackler, J. Leggett, G.M. McColvin, S. Simpson, M. Smith, F. Star, R.W. Vanstone (Eds.), Proceedings of the Sixth International Charles Parsons Turbine Conference, 16–18 September 2003, Trinity College Dublin, Ireland, 2003, pp. 261–279. [4] C. Middleton, E. Metcalfe, IMechE Proceedings, London, UK, 1990 (Paper C386/027). [5] Shibli, Proceedings of the Swansea Creep Conference, University of Swansea and EPRI, Swansea, UK, 2001. [6] J.A. Francis, W. Mazur, H.K.D.H. Bhadeshia, Mater. Sci. Technol. 22 (12) (2006) 1387–1395. [7] H. Cerjak, E. Letofsky, in: R. Viswanathan, J. Nutting (Eds.), Advanced Heat Resistance Steels for Power Generation, 1998, pp. 611–621. [8] J.M. Brear, A. Fairman, C.J. Middleton, L. Polding, Key Eng. Mater. 171–174 (3) (2000) 5–42. [9] C.J. Middleton, J.M. Brear, R. Munson, R. Vishwanathan, in: R. Vishwanathan, W.T. Bakker, J.D. Parker (Eds.), Proceedings of the 3rd Conference on Advances in Materials Technology for Fossil Power Plant, The Inst. Mater., London, 2001, pp. 69–78. [10] E. Letofsky, H. Cerjak, I. Papst, P. Warbichler, in: R. Vishwanathan, W.T. Bakker, J.D. Parker (Eds.), Proceedings of the 3rd Conference on Advances in Materials Technology for Fossil Power Plant, The Inst. Mater., London, 2001, pp. 133–142. [11] F. Masuyama, M. Matsui, N. Komai, Key Eng. Mater. 171–174 (2000) 99–108. [12] M. Matsui, M. Tabuchi, T. Watanabe, K. Kubo, J. Kinugawa, F. Abe, ISIJ Int. 41 (2001) S126–S130. [13] M. Tabuchi, T. Watanabe, K. Kubo, M. Matsui, J. Kinugawa, F. Abe, Int. J. Press. Vessels Piping 78 (2001) 779–784. [14] M. Tabuchi, M. Matsui, T. Watanabe, H. Hongo, K. Kubo, F. Abe, Mater. Sci. Res. Int. 9 (2003) 23–28. [15] F. Vivier, Fluage a’ 500 ◦ C d’ un joint sounde’ d’ un acier 9Cr–1Mo modifie’. Evolution de la microstructure & Comportement me’canique, PhD thesis, 2009, Ecole des Mines de Paris, France. [16] D. Jandova, J. Kasl, V. Kanta, Proceedings of the 2nd International ECCC Conference, Empa, CH-8600 Dübendorf Akademie, Überlandstrasse 129, April, 2009. [17] V. Gaffard, Experimental study and modeling of high temperature creep flow and damage behaviour of 9Cr–1Mo–NbV steel weldments, PhD thesis, 2004, Ecole des Mines de Paris, France. [18] T. Watanbe, M. Tabuchi, M. Yamazaki, H. Hongo, T. Tanabe, Int. J. Press. Vessels Piping 83 (2006) 63–71. [19] ASTM E606-92, Standard recommended practice for constant-amplitude lowcycle fatigue testing, Annual Book of ASTM Standards, 1994, 03.01, pp. 522–536. [20] V. Shankar, M. Valsan, K. Bhanu Sankara Rao, R. Kannan, S.L. Mannan, S.D. Pathak, Mater. Sci. Eng. A 437 (2) (2006) 413–422. [21] K.S. Chandravathi, K. Laha, K.B.S. Rao, S.L. Mannan, Mater. Sci. Technol. 17 (2001) 559–565. [22] D.P. Edmonds, B. Chew, Effect of chemical composition on weld microstructure in some commercial 9Cr–1Mo steel boiler tubes. Report TPRD/M/1258/N82. Central Electricity Generating Board, Marchwood, 1982. [23] K. Bhanu Sankara Rao, M. Valsan, R. Sandhya, S.K. Ray, S.L. Mannan, P. Rodriguez, Int. J. Fatigue 7 (3) (1985) 141–147.
V. Shankar et al. / Materials Science and Engineering A 528 (2011) 8428–8437 [24] S. Kim, J.R. Weertman, Metall. Trans. A 19 (1988) 999–1007. [25] B. Fournier, M. Sauzay, A. Renault, F. Barcelo, A. Pineau, J. Nucl. Mater. 386–388 (386) (2009) 71–80. [26] H. Li, A. Nishimura, T. Nagasaka, T. Muroga, Fusion Eng. Des. 81 (2006) 2907–2912. [27] A.F. Armas, C. Petersen, R. Schmitt, M. Avalos, I. Alvarez-Armas, J. Nucl. Mater. 307–311 (307) (2002) 509–513. [28] S.T. Kimmins, M.C. Coleman, D.J. Smith, in: B. Wilshire, R.W. Evans (Eds.), Proceedings of the Fifth International Conference on Creep and Fatigue of Engineering Materials and Structure, Swansea, 1993. [29] D.C. Lord, L.F. Coffin Jr., Metall. Trans. 4 (1973) 1647–1654. [30] W.B. Jones, in: A.K. Khare (Ed.), Proceedings of the ASM International Conference on Production, Fabrication, Properties and Application of Ferritic Steels for High Temperature Applications, Warren, PA, 1981, pp. 221–235.
8437
[31] P. Marmy, T. Kruml, J. Nucl. Mater. 377 (2008) 52–58. [32] B. Fournier, M. Sauzay, C. Caës, M. Noblecourt, M. Mottot, A. Bougault, V. Rabeau, J. Man, O. Gillia, P. Lemoine, A. Pineau, Int. J. Fatigue 30 (October–November (10–11)) (2008) 1797–1812. [33] J.S. Dubey, H. Chilukuru, J.K. Chakravartty, M. Schwienheer, A. Scholz, W. Blum, Mater. Sci. Eng. A 406 (October (1–2)) (2005) 152–159. [34] B. Fournier, M. Sauzay, C. Caës, M. Noblecourt, M. Mottot, A. Bougault, V. Rabeau, A. Pineau, Int. J. Fatigue 30 (April (4)) (2008) 649–662. [35] B. Fournier, M. Sauzay, C. Caës, M. Noblecourt, M. Mottot, A. Bougault, V. Rabeau, A. Pineau, Int. J. Fatigue 30 (April (4)) (2008) 663–676. [36] S. Goyal, K. Laha, K.S. Chandravathi, P. Parameswaran, M.D. Mathew, Philos. Magazine 91 (23) (2011) 3128–3154. [37] I. Nonaka, S. Torihata, S. Kihara, H. Umaki, Mater. High Temp. 15 (2) (1998) 69–73.