FR0202659 Reheat cracking in austenitic stainless steels 1 1 2 2 Q.Auzoux , L.AlIais , A.Pineau , A.F.Gourgues 1 CEA Saclay DEN/DMN/SRMA Bat 455 91191Gif sur Yvette France 2 Centre des Matériaux Pierre-Marie Fourt UMR CNRS 7633 B.P. 87 91003 Evry France ABSTRACT: Intergranular cracking can occur in heat-affected zones (HAZs) of austenitic stainless steel welded joints when reheated in the temperature range from 500 to 700°C. At this temperature, residual stresses due to welding relax by creep flow. IIAZ may not sustain this small strain if its microstructure has been sufficiently altered during welding. In order to precise which particular microstructure alteration causes such an intergranular embrittlement, type 316L(N) HAZs were examined by transmission electron microscopy. A marked increase in the dislocation density, due to plastic strain during the welding process, was revealed, which caused an increase in Vickers hardness. Type 316L(N) HAZ were then simulated by the following thermal-mechanical process: annealing treatment and work hardening (pre-strain). Creep rupture tests on smooth specimens were also carried out at 600°C on both base metal and simulated HAZ. Pre-straining increased creep strength but reduced ductility. Slow strain rate tests on CT specimens confirmed this trend as well as did relaxation tests on CT specimens, which led to intergranular crack propagation in the pre- strained material only. Metallography and fractography showed no qualitative difference between base metal and HAZs in the creep cavitation around intergranular carbides. Although quantitative study of damage development is not achieved yet, experiments suggest that uniaxial creep strain smaller than one percent could lead to cavity nucleation when the material is pre-strained. Pre-strain as well as stress triaxiality reduce therefore creep ductility and enhance the reheat cracking risk. INTRODUCTION Plant experience and research studies on reheat cracking in austenitic stainless steels have been reviewed by Dhooge [1]. Not only stabilized stainless steels but also type 316 stainless steel appeared to be prone to reheat cracking [2]. This phenomenon is consistent with the following result of Chabaud-Reytier [ 3 J : the reheat cracking sensitivity of a 321 stainless steel was found to depend drastically on the cold work level rather than on the precipitation state. In an attempt to confirm this effect of pre-strain on reheat cracking, we chose a low carbon, high nitrogen 316L(N) stainless steel less sensitive to intragranular precipitation. HAZs of actual welded joints were examined by transmission electron microscopy (TEM) and a
1 CEA Saclay DEN/DMN/SRMA Bat 455 91191Gif sur Yvette France 2 Centre des Matériaux Pierre-Marie Fourt UMR CNRS 7633 B.P. 87 91003 Evry France
ABSTRACT: Intergranular cracking can occur in heat-affected zones (HAZs) of austenitic stainless steel welded joints when reheated in the temperature range from 500 to 700°C. At this temperature, residual stresses due to welding relax by creep flow. IIAZ may not sustain this small strain if its microstructure has been sufficiently altered during welding. In order to precise which particular microstructure alteration causes such an intergranular embrittlement, type 316L(N) HAZs were examined by transmission electron microscopy. A marked increase in the dislocation density, due to plastic strain during the welding process, was revealed, which caused an increase in Vickers hardness. Type 316L(N) HAZ were then simulated by the following thermal-mechanical process: annealing treatment and work hardening (pre-strain). Creep rupture tests on smooth specimens were also carried out at 600°C on both base metal and simulated HAZ. Pre-straining increased creep strength but reduced ductility. Slow strain rate tests on CT specimens confirmed this trend as well as did relaxation tests on CT specimens, which led to intergranular crack propagation in the pre-strained material only. Metallography and fractography showed no qualitative difference between base metal and HAZs in the creep cavitation around intergranular carbides. Although quantitative study of damage development is not achieved yet, experiments suggest that uniaxial creep strain smaller than one percent could lead to cavity nucleation when the material is pre-strained. Pre-strain as well as stress triaxiality reduce therefore creep ductility and enhance the reheat cracking risk.
INTRODUCTION
Plant experience and research studies on reheat cracking in austenitic stainless steels have been reviewed by Dhooge [1]. Not only stabilized stainless steels but also type 316 stainless steel appeared to be prone to reheat cracking [2]. This phenomenon is consistent with the following result of Chabaud-Reytier [ 3 J : the reheat cracking sensitivity of a 321 stainless steel was found to depend drastically on the cold work level rather than on the precipitation state. In an attempt to confirm this effect of pre-strain on reheat cracking, we chose a low carbon, high nitrogen 316L(N) stainless steel less sensitive to intragranular precipitation. HAZs of actual welded joints were examined by transmission electron microscopy (TEM) and a
thermal-mechanical treatment of HAZ simulation was chosen to reproduce a typical HAZ microstructure. The susceptibility to reheat cracking of this simulated HAZ was then estimated using a relaxation test on compact-tension (CT) specimen as described in [4]. Tensile and creep tests were also performed in order to compare the mechanical behaviour and fracture properties of base metal and simulated HAZ, respectively.
MATERIALS AND EXPERIMENTAL PROCEDURES
Materials Type 316L(N) stainless steel was supplied by A vesta Sheffield. The chemical composition is (wt%): C: 0.026, N: 0.069, Cr: 17.3, Ni: 12.1, Mo: 2.54, Mn: 1.74, Si: 0.31, Cu: 0.29, S: 0.001, P: 0.025, Co: 0.089, B: 0.004. The as received, 25mm thick plates were annealed at 1170°C during 30 minutes and gas cooled. During this heat-treatment, the average grain size increased from 35 to 60 pm. Part of this annealed material was then rolled between 400°C and 600°C to reduce the thickness by 15%. The Vickers hardness of the annealed material and of the pre-strained material is 140 HV30 and 230 HV30 respectively. Tensile properties of both materials at 600°C are reported in Table L A dynamic strain-ageing effect was noticed during these tests, except for strain rates smaller than 10"4 s"1. A manual metal arc cross-weld was manufactured using the as-received plate. Heat input was 0.5 kJ/cm. Vickers hardness measurements in the vicinity of this weld are plotted in figure 1. The HAZ is much harder than the base metal and a little softer than pre-strained material. The dislocation structure of both real HAZ and the pre-strained material is composed of cells (figure 2). The dislocation density is much higher than in the base metal. Fewr sub-grains were also observed in real HAZ. Before ageing, a small amount of intergranular M23C6, but no intragranular precipitation, was evidenced in both real HAZ and pre-strained materials. Hardening of the HAZ is therefore thought to be caused by the increase in dislocation density due to the deformation occurring during welding process. Pre-strained material is representative from the hardest part of the HAZ. Both hardness measurements and microstructure observations are consistent with those reported by Etienne [5].
Experimental procedures Creep and tensile tests were performed on smooth specimens cut perpendicularly to the rolling direction. Diameter was 4 mm, gauge length
20 mm and total length 40 mm. Temperature, controlled by three thermocouples attached to the specimen was 600°C±3°C. Slow strain rate tests and relaxation tests were also realised on pre-cracked CT specimens. The specimens were 10 mm thick (side-grooved to 8 mm) and 40 mm wide. The load was applied transversally to the rolling direction, and the crack propagated parallel to the rolling direction. A potential-drop technique was used to evaluate the crack growth. Pre-cracking was realised such as a<)AV = 0.5. Opening displacement was recorded through a capacitive transducer attached to the top and bottom of the specimen along the loading line. Some of the tests were performed under high vacuum, with pressure lower than 4.10'4Pa. During slow strain rate tests, the opening displacement rate imposed was 5.10"6 mm/s. The test was interrupted after the load had reached a maximum value. During relaxation tests the specimen was first quickly loaded (4.10" mm/s) up to a pre-set value of the load. Then, the opening displacement was monitored to remain constant (± 0.2pm) and the progressive decrease in the load was recorded.
Figure 1 : Vickers hardness in the vicinity of the cross weld.
i . f § |
• 1/4 thickness
o 1/2 thickness
i . § • ç - I
Figure 2: TEM bright field micrographs of (left) real HAZ. 200 pm far from the fusion line, and (right) simulated HAZ (pre-strained material).
RESULTS AND DISCUSSION
Creep tests on smooth specimens revealed marked differences between both materials. Plastic elongation on loading occurred only for the annealed material (5 to 9 %) because of the difference in flow stress between the materials. Figure 3 shows that for the same true stress (nominal stress corrected by section reduction on loading), pre-strain caused at least a 400-fold decrease in the minimum creep rate and a 20-fold increase in the time to failure. The strengthening effect of room-temperature pre-strain (cold work) had already been observed in type 316 and 304 stainless steels [6-10]. Our results show that pre-straining between 400°C and 600°C has a similar effect. Both the increase in dislocation density and the strain-induced precipitation of intragranular carbides have been quoted as main explanation of this strengthening effect. TEM examinations on crept materials will be performed to check this point. However, the low carbon, high nitrogen content of the present heat leads us to assume a weak intragranular precipitation before 100 h at 600°C [11] and therefore to presume that the increase in dislocation density is the major reason for the observed strengthening. Moreover, the prc-strain effect decreases when time to rupture increases, which is more coherent with a recovery-controlled creep, rather than a particle-controlled creep. Ductility is known to decrease with strain rate in the range 10'4 s'1 to 10"s s"1
because of the progressive change in the fracture mechanism from transgranular ductile to more brittle intcrgranular. Observation of the
fracture surfaces and of polished cross-sections confirmed this trend as plotted in figure 4. However, we would like to emphasize that, even if a large amount of intergranular micro-cracks are observed in the specimen, final fracture occurs by localized shear and dimple growth, so that fracture surfaces seldom showed grain boundary facets. Both annealed and pre-strained materials experienced such a transition in the rupture mechanism, but figure 4 indicates that for the same creep rate, pre-strained material is more brittle than annealed material. The Monkman-Grant product: £s.tR is approximately equal to 1 1% for annealed material but only to 1.3% for pre-strained material. The elongation to failure also decreases from the range 15-40% for annealed material to the range 7-15% for pre-strained material. This ductility drop explains why the strengthening effect of pre-strain is not as high on creep life as on creep flow.
400
340
•320
300
280
260
240
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200
1 • Pre-strained
: o Annealed
1 • Pre-strained
: o Annealed t
I •
c
E
\
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360 « a 340 3?0 v> 300
v> ?R0 o 260
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I v 1 • Pre-strained
o Annealed \ 1
• Pre-strained
o Annealed
Î 1.E-10 1.E-9 1.E-8 1.E-7 1.E-6
Min imum crccp rate (s-1) 1.E-5 10 100 1000
Time to rupture (hours) 10000
Figure 3: Creep tests on smooth specimens tested at 600°C 1 T
Abundant intergranular microcracks
O O • o
o
Dimples & Shear °
•
o
•
Scarso intergranular microcracks
Dimples & Shear
r • Dimples & Shear
• • Pre-strained
o Annealed i ••' 1 " " " * * 1 1
: Creep tests
Partly intergranular fracture
Tensile tests
Trans-granular fracture
[-10 1.E-8 1.E-6 1.E-4 Min imum creep rate or strain rate s-1 Minimum creep rate or strain rate (s-1)
Figure 4: Creep and tensile tests on smooth specimens at 600°C
Pre-strain embrittlement (PSH) was also noticed on CT specimens as illustrated in figure 5 for slow strain rate tests under vacuum. Similar tests carried out in air led to similar results, indicating that intergranular cracking was not controlled by oxidation phenomena. The most relevant test that reveals PSE is the relaxation test on CT specimens. As reported in table 2, annealed material never led to crack growth, even after lOOOh, in opposition to pre-strained material, which started to crack after only a few hours. Figure 6 shows the intergranular feature of such a crack growth.
PSE may be caused by an increase in the intergranular cavity nucleation linked to local stresses around intergranular particles due to pre-strain, as described in cold worked Nimonic 80 [12]. However, pre-strained 316 stainless steel differs from the latter. After testing, creep cavities around intergranular carbides (=0.15um) or at triple junctions were observed near the main crack of the CT specimens, but not away from it, even when sections were carefully polished and observed by field emission gun scanning electron microscope. Therefore, intergranular cavities are likely to continuously nucleate with creep strain as a result of local stress concentration due to deformation incompatibility [13]. The effect of pre-strain could be to increase these local stresses for a given creep strain. Reheat-cracking risk is enhanced by residual stress triaxiality and slow strain rate [14,15]. Our results tend to prove that plastic strain during welding also has a deleterious effect.
Opening displacement (mm)
Figure 5: Slow strain rate test on CT specimens tested under vacuum at 600°C, opening rate = 5.10"6 mm/s. Intergranular crack propagation Aa = 11.8 mm and 2.7 mm for pre-straincd and annealed material, respectively
Figure 6: SEM micrographs of the fracture surface of a CT specimen relaxation tested in air at 600°C, pre-strained material
CONCLUSION
Type 316L(N) HAZs show a marked increase in dislocation density due to plastic strain occurring during welding, which also causes an increase in yield stress. A similar microstructure was obtained using a work hardening process around 500°C. Creep tests on smooth specimens revealed a beneficial effect of pre-strain on creep resistance but also a drop in ductility. This pre-strain embrittlement (PSE) was also noticed during slow strain rate tests and relaxation tests on CT specimens. PSE may be the consequence of an enhancement of intergranular cavity nucleation due to the increase in local stresses by pre-strain. Pre-strain during welding should therefore be as small as possible to limit the risk of reheat cracking in austenitic stainless steels.
REFERENCES
1. Dhooge A. (1998) Welding in the world, 41, pp 206-219. 2. Coleman M.C., Miller DA. , Stevens R A . (1998) Integrity of High-
Temperature Welds. Proceedings, Int. Conf. Nottingham, pp. 169-179. 3. Chabaud-Rcytier M. (1999) Ph.D. Thesis, Ecole Nationale Supérieure
des Mines de Paris, France. 4. Poquillon D., Chabaud-Reytier M., Allais L., Pineau A. (2001) Mat. at
High Temp., 18, 2, pp.81-90. 5. Etienne C.F., van Rossum O., Roode F. (1980) International conference
on engineering aspects of creep, Sheffield, 2, pp. 113-121. 6. Samuel K.G., Bhanu Sankara Rao K., Mannan S.L., Radhakrishnan
V.M. (1996) Trans. Indian Inst. Met., 42, 4, pp. 407-412. 7. Dyson B.F., Loveday M.S. (1980) International conference on
engineering aspects of creep, Sheffield, 1, pp. 61-66. 8. Wilshire B., Willis M.R. (2001) 10th. International conference on creep
and fracture of engineering materials and structures, Prague, pp. 6-15. 9. Ajaja O., Ardell A.J. (1977) Script. Met., 11, pp. 1089-1093. 10. Bernard L., Campo E., Quaranta S. (1981) Mechanical behaviour and
nuclear applications of stainless steel at elevated temperatures, Conf. Varese 1981, pp. 88-92.
11. Weiss B., Stickler R. (1972) Met. Trans. 3, pp. 851-866. 12. Dyson B.F., Rodgers M.J. (1974) Met. Sci., 8, pp. 261-266 13. Dyson B.F. (1983) Script. Met. 17, pp. 31-37. 14. Spindler M.W. (2001) HTMTC Meeting on Multiaxial Creep Testing
and Interpretation, IOM London, to be published in Fatigue Fract. Engng. Mater. Struct.
15. Hales R. (1994) Fatigue Fract. Engng. Mater. Struct. 17. 5, pp. 579-591.
ACKNOWLEDGEMENTS
The authors would like to thank I. Tournié, C. Caës, from CEA and R. Lociccro from CdM for performing mechanical tests. Thanks also to I. Monnet and T. van den Berghe from CEA for TEM examinations.
