- Email: [email protected]
Embrittlement of laser surface-annealed 17-4 PH stainless steel
Materials Science and Engineering A311 (2001) 64 – 73 www.elsevier.com/locate/msea
Embrittlement of laser surface-annealed 17-4 PH stainless steel L.W. Tsay *, T.Y. Yang, M.C. Young Institute of Materials Engineering, National Taiwan Ocean Uni6ersity, 2 Pei-Ning Road, Keelung 202, Taiwan, ROC Received 10 October 2000; received in revised form 18 December 2000
Abstract Stress corrosion cracking and fatigue crack growth behavior were determined in 17-4 PH stainless steel under various metallurgical conditions, including the H900 (482°C/1 h), H1025 (552°C/4 h) aging and laser surface annealing treatments. Peak-aged (H900) specimens locally irradiated by laser beam consisted of a portion of composite region (CR), in which comprised of soft laser-annealed (LA) zones on the outer surfaces and the hard base metal in between. Slow extension rate tensile tests were performed at room temperature in a saturated H2S solution to evaluate the hydrogen embrittlement susceptibility of various specimens. H900 specimens show an obvious improvement in impact toughness after irradiating by laser. Regardless of testing environments, H900 specimens exhibited better tensile properties than the other specimens. Experimental results also indicated that H900 specimens had the highest fatigue crack growth rates among the specimens, particularly at low stress intensity factor range. The retardation of crack growth in the region ahead of the CR in the LA specimens was rather pronounced. For compact tension specimens tested in gaseous hydrogen, enhanced crack growth was correlated with quasi-cleavage fracture in contrast to transgranular fatigue fracture in air. © 2001 Elsevier Science B.V. All rights reserved. Keywords: Fatigue crack growth; Hydrogen embrittlement; Composite region; Laser annealing treatment
1. Introduction 17-4 PH (AISI 630) steel is a precipitation-hardening martensitic stainless steel, which behaves high strength, high toughness and moderate corrosion resistance [1,2]. This steel can be used in differently aged conditions and is widely employed in various industries. The microstructure in the solution-annealed (SA) condition comprises a mixture of a martensite and 5 – 10 volume percent d ferrite, the latter elongated in the prior working direction [3]. A rapid increase in the hardness of the alloy during aging is owing to the initial formation of coherent copper-rich cluster in the peak-aged condition [4]. At temperature up to 600°C, overaging occurs resulting in the formation of incoherent o phase precipitates together with the transformation of a significant amount of martensite to austensite along martensite lath boundaries [5]. The microstructural changes not only alter the mechanical properties but also affect the corrosion resistance [3]. In general, the steel in the * Corresponding author. Fax: + 886-2-24625324. E-mail address: [email protected] (L.W. Tsay).
overaged condition has a higher resistance to sulfide stress corrosion cracking (SSCC) because of its lower hardness [6]. Both KIC and KISCC (in 3.5% aqueous sodium chloride solution) decreased inversely with strength level for high strength precipitation hardening stainless steels [2]. However, the precipitation of incoherent o phase in overaged specimens caused the depletion of Cu and Ni in the matrix, resulting in the enhanced electrochemical dissolution of the martensite matrix [3]. The use of laser irradiation in material processing has been increased in recent years. If the steel has a suitable carbon content after laser treatment, the near-surface region heated to above the austenizing temperature will be transformed to hard martensite [7]. The short thermal cycle of the laser processing can also result in the formation of softened zones in the overlapped areas after multiple laser passes [8]. The hardened welds could also be tempered effectively by controlling the power and irradiating area of the laser beam [9]. Depending on the thermal cycles at various locations, the fine precipitates in the surface region of the peak-aged specimen after laser irradiation may be dissolved into
0921-5093/01/$ - see front matter © 2001 Elsevier Science B.V. All rights reserved. PII: S 0 9 2 1 - 5 0 9 3 ( 0 1 ) 0 0 9 4 1 - 8
L.W. Tsay et al. / Materials Science and Engineering A311 (2001) 64–73
the matrix or coarsened. Hence, a specific depth of surface layer with lower hardness than the interior hard base metal (BM) is formed, and the soft surface layer is defined as the laser-annealed zone (LAZ). Locally symmetrical irradiation on the top and bottom surfaces of peak-aged specimens produced a composite region (CR), in which consisted of soft LAZs on the outer surfaces and untransformed BM sandwiched between them. In general, high strength steels are sensitive to hydrogen embrittlement and surface notches. It was reported [10] that both hydrogen embrittlement and anodic dissolution mechanisms are responsible for the stress corrosion cracking (SCC) of 17-4 PH steel in an acidic NaCl solution. Under applied anodic potentials, the 17-4 PH stainless steel in the SA state was found to be more resistant to SCC [11]. It is well known that an important factor affecting SSCC susceptibility is the strength level of the material, which is directly related to the material hardness. As defined in NACE standard MR0175-88, the maximum allowable hardness for 17-4 PH stainless steel is HRC33 to resist SSCC [6]. 17-4 PH stainless steel is extensively used in the power industry. In geothermal power application, structural components may suffer from hydrogen embrittlement, owing to the presence of H2S species in the environment. Laser irradiation on the peak-aged 17-4 PH stainless steel leads to the formation of soft surface layer in the SA condition with high ductility/toughness. The combination of soft LAZ on the surface and interior strong BM may not only possess adequate strength/ductility, but also reduce the susceptibility to SSCC and improved resistance to crack growth of the alloy. This study was undertaken with the objective of characterizing the influence of aging condition and hydrogen charging on the tensile and fatigue crack growth behavior in 17-4 PH stainless steel aged at various conditions. It is reported that SSCC was a manifestation of hydrogen embrittlement [12,13]. Slow extension rate tensile tests were usually performed to evaluate the susceptibility to hydrogen embrittlement of various specimens in an H2S-saturated NACE solution [14,15]. Impact energy of the specimens measured at room temperature was used to assess the influence of laser irradiation on the improved resistance to unstable brittle fracture, and the results were compared with aged specimens. It is found that few research works have been performed to study the fatigue crack growth behavior in 17-4 PH stainless steel aged at various conditions. Compact-tension (CT) specimens tested in air and gaseous hydrogen were used to determine the fatigue crack growth rates of various specimens, and the effect of hydrogen embrittlement on accelerating crack growth was evaluated. Two types of fatigue tests, constant load and constant DK were conducted in this
65
work. Fractographic observations were performed to correlate the microstructures and changes in properties with fracture features.
2. Materials and experimental procedures The chemical composition in wt.% of the 17-4 PH stainless steel used in this investigation in the plate form of 5.6 mm thick was 15.71Cr, 4.35Ni, 3.46Cu, 0.04C, 0.7Si, 0.38Mn, 0.025P, 0.013S and balance Fe. The round bar of 15 mm diameter was composed of (wt.%) 16.60Cr, 4.26Ni, 3.50Cu, 0.046C, 0.52Si, 0.6Mn, 0.03P, 0.043S and balance Fe. The alloy was SA at 1038°C for 1 h and then air-cooled to room temperature. The aging treatments were performed either at 482°C for 1 h (H900) or 552°C for 4 h (H1025). Some of the H900 specimens were subjected to a laser-annealed (LA) treatment. In this investigation, laser beam integrators (6× 6 mm2) and rectangular beams (6×25 mm2) were used to provide uniform distribution of energy over a relatively large area, resulting in the formation of a more homogeneously annealed region without surface-melting of the specimen. A Rofin –Sinar 5 KW CO2 laser connected with a CNC table was utilized for laser treatment. Steel plates aged at 482°C were processed by laser at 2000 W and 800 mm min − 1 scan rate and the round specimens were laser-treated at 600 W with 2 turns s − 1 for 5 s. The hardness of the steel plate in various conditions, including SA, as well as H900 and H1025 aging treatment, was measured by Rockwell hardness tester (HRC). After laser treatment, some work pieces were sectioned transverse to the laser scan direction (LSD), then polished and etched in a Vilella’s reagent. Microhardness measurements across the LAZ was taken via a Vickers microhardness tester. Fig. 1 shows the schematic dimensions of tensile and Charpy impact specimens employed in this study. Two types of tensile specimens with a gauge length of 25 mm are shown in Fig. 1. The tensile properties of various specimens tested in air were determined at an ordinary strain rate of 5×10 − 3 s − 1. In order to evaluate the hydrogen embrittlement susceptibility of various specimens, slow extension rate tensile tests were carried out at room temperature in an H2S-saturated solution with a constant strain rate of 5× 10 − 6 s − 1. The test solution prepared according to the NACE standard (TM-01/7786) was purged with N2 for at least 1 h to get rid off O2 before testing and then H2S was bubbled continuously to ensure that the solution was saturated with H2S. The results represent the average of at least three specimens for each testing condition. Charpy impact specimens (Fig. 1(c)) are made according to the specifications of ASTM E23 (subsize, 5 mm thick) fractured at room temperature.
66
L.W. Tsay et al. / Materials Science and Engineering A311 (2001) 64–73
The fatigue crack growth tests were performed at room temperature in a computerized hydraulic testing machine (MTS 810). CT specimens are used and the experimental procedures meet the standard of ASTM E647-91 specification. The geometry of the specimens used in this study is shown in Fig. 2, in which the crack growth direction is normal to the LSD. The testing software (759.40 crack growth test) developed by MTS can offer a variety of testing requirements including constant load, constant DK and K-controlled step testing conditions. In order to magnify the differences in fatigue crack growth rates in various regions of a LA specimen, constant DK tests were performed at a stress ratio of 0.1. In addition to the CT specimen that were tested in air, the effect of hydrogen embrittlement on
Fig. 2. Schematic diagram showing the dimensions of CT specimen.
accelerating the crack growth was investigated by testing the CT specimens inside a stainless steel chamber with 2 atm. hydrogen gas. In such a condition, a capital ‘‘H’’ is added to the last symbol of the designated specimen, e.g. H900-H represents the hydrogen-charged H900 specimen. Fractographic observations of the fractured specimens were examined by a Hitachi 4100 scanning electron microscope (SEM), with emphasis paid to the change in fracture modes. Residual stresses measurement (ASTM E837-92) was performed by a modified hole-drilling strain gauge method. Three-element strain gauge rosettes (Measurements Group, CEA-06-062UL-120) were applied onto the CT specimen at the centerline of the LAZ and 6 mm away from the centerline (in the BM). A hole of 1.6 mm diameter was drilled by an air turbine hole-drilled machine in the center of the rosette. Relaxed strains in terms of residual stresses were measured by a strain indicator.
3. Results and discussion
3.1. Hardness measurements and metallographic obser6ations
Fig. 1. Schematic diagrams showing the dimensions of (a) round tensile specimen, (b) rectangular tensile specimen and (c) Charpy impact specimen; dimension in mm.
Both the Rockwell and Vickers hardness testers were used to measure the hardness of 17-4 PH steel plates aged in various conditions. The initial hardness is of about HRC32.1 (HV370) for the specimen in the SA condition. After aging at 482°C for 1 h (H900), the precipitation reaction resulted in an increase in hardness to HRC44.5 (HV465). Aging at a higher tempera-
L.W. Tsay et al. / Materials Science and Engineering A311 (2001) 64–73
ture, e.g., 552°C (H1025), caused the hardness to decline to HRC37.2 (HV389). Fig. 3 displays schematic diagram of microhardness distribution in LA specimens measured at the center of the LA region for the round and rectangular tensile specimens. A decrease in surface hardness to HV370 similar to the hardness of the SA specimen was noted. It suggested that the near-surface
67
Table 1 Tensile properties of 17-4 PH steel in various conditions tested in air and H2S solution for round tensile specimensb Materials Properties/Condition
UTS (MPa) Elongation (%) Reduction in area (%) Fracture location
H900
H900+LA
In Aira
In H2Sb
In Air
In H2S
1273 19.3 30
1157 7.5 15.2
1181 10.3 27.1
1099 6.5 15.0
–
–
CR
CR
a UTS – ultimate tensile strength, CR – composite region. In Air – tested in laboratory air with the strain rate of 5×10−3 s−1. b In H2S – tested in an H2S-saturated solution with the strain rate of 5×10−6 s−1.
Fig. 3. Schematic diagram showing microhardness distribution in the rectangular and round LA specimens.
region after laser irradiation should be transformed back to the SA condition. In addition, the variation in microhardness distribution in the LA specimen could be attributed to the complex matrix/precipitate reactions resulting in a continuous phase transformation within this narrow heat-affected region. Fig. 4 reveals the metallographs of the various specimens. The microstructures of the H900 specimen consisted of a mixture of martensite and elongated d ferrite oriented along the rolling direction (Fig. 4(a)). The SEM/EDX spectrum revealed that the Cr concentration in the d ferrite was slightly higher than that in the matrix. It has been found that the near-surface region in the LAZ contained less stringer d ferrite (upper portion of Fig. 4(b)). It should be due to the high temperature thermal cycle associated with laser treatment resulting in the dissolution of d ferrite into the matrix. Hence, elongated d ferrite was refined at the near-surface region. It had been pointed out previously [3] that the compositional difference in Cr could account for the localized attack at the d ferrite/matrix interface. Thus, improvement in corrosion resistance of the LAZ could be anticipated, owing to the dissolution of d ferrite as well as the reduction of d ferrite/matrix interface.
3.2. Tensile properties
Fig. 4. Metallographs showing (a) the elongated d ferrite in the steel plate and (b) refined d ferrite in the near-surface region of an LA specimen. Arrows indicating the stringer d ferrite.
Table 1 and Table 2 present the tensile properties of 17-4 PH stainless steel aged in various conditions tested in air and H2S solution for round and rectangular tensile specimens, respectively. In case of the round tensile specimen annealed by laser, a specific depth of soft surface layer (i.e. LAZ) enclosing the interior hard BM is shown in Fig. 1(a). The CR of the round tensile specimen included the LAZ at the periphery and interior BM. As shown in Fig. 1(b), semi-elliptical LAZs were located on the top and bottom surfaces and BM sandwiched between them in rectangular tensile speci-
68
L.W. Tsay et al. / Materials Science and Engineering A311 (2001) 64–73
Table 2 Tensile properties of 17-4 PH steel in various conditions tested in air and H2S solution for rectangular tensile specimens Materials Properties/Condition
UTS (MPa) Elongation (%) Reduction in area (%) Fracture location
H900
H1025
H900+LA
In Air
In H2S
In Air
In H2S
In Air
In H2S
1356 19.2 50.1 –
948 3.6 2.2 –
1087 21.7 57 –
822 4.2 3.6 –
1234 14.3 49.2 CR
822 4.1 3.0 CR
mens. In this case, the BM in the CR was not completely enclosed by the LAZs. It was noted that a slightly different degree of embrittlement was found for those two types of CR testing in a corrosive environment. The sandwiched BM in the CR of a rectangular tensile specimen would directly immerse into H2S solution, in contrast to that of round tensile specimens, the BM was surrounded by the LAZ. For the round tensile specimen tested in air, the H900 specimen had slightly better tensile properties than the LA specimen. When tested in an H2S-saturated solution, both the H900 and LA specimen showed a slight decrease in strength but greater reduction in ductility as compared to that tested in air. The fracture location of the LA specimen, regardless of testing environment, was in the CR. As given in Table 2, rectangular tensile specimens show the same change of trend in tensile properties as round tensile specimens. H900 specimens possessed the best tensile properties among specimens tested in air, and H1025 specimens had lower strength than the other specimens. When tested in an H2S-saturated solution, H900 specimens still exhibited the highest strength among the various specimens. Meanwhile, the results also indicated that specimen’s geometry had a great influence on the tensile strength of the specimens tested in an H2S-saturated solution. It was known that the fracture stress of the material in hydrogen decreased as the notch severity was increased [16]. A more obvious degradation in tensile strength for a rectangular specimen could be attributed to the more inhomogeneous distribution of tensile stress over the cross-section of the specimen. Subsequently, all the specimens were immersed in the saturated H2S solution and the surface cracks were inspected after a period of time. The optical micrographs revealing the surface morphology of immersed specimens are shown in Fig. 5. It was found that surface cracks were initiated in the H1025 specimen after being immersed in an H2S-saturated solution for 7 h (Fig. 5(a)) but no detectable surface cracks were found in other specimens. Charging the specimens in H2S for 20 h yielded the same surface defects on the specimens (LA, SA and H900 specimens) i.e., corrosion
pits were formed and eventually coalesced into cracks (Fig. 5(b)). Such results demonstrated that SA, LA and H900 specimens had the similar ability to resist hydrogen-induced cracking, meanwhile, H1025 specimen was more sensitive to hydrogen embrittlement. SSCC susceptibility in various region of a weld has been found to vary with the applied stress level [14]. It was reported [14] that SSCC could occur at the relatively soft region under a high stress level. In this work, the immersion tests showed that the H900 specimen revealed similar susceptibility to hydrogen-induced cracking as the LAZs within the testing time interval. It was pointed out previously that SA specimen was
Fig. 5. Optical micrographs revealing (a) surface cracks in the H1025 specimen immersed in an H2S-saturated solution for 7 h, and (b) corrosion pits in the H900 specimen immersed in an H2S-saturated solution for 20 h.
L.W. Tsay et al. / Materials Science and Engineering A311 (2001) 64–73
69
decrease in tensile strength and ductility of LA specimens. It has also been reported that d ferrite in precipitation age-hardened stainless steels will increase the susceptibility to hydrogen embrittlement [17]. When the testing strain rate was decreased or the severe notch was introduced, the CR might show lower susceptibility to hydrogen embrittlement and better resistance to SSCC than the hard BM, which was related to the less stringer d ferrite and lower surface hardness.
3.3. Charpy impact test
Fig. 6. Macroscopic photograph of tensile fractured specimens tested in an H2S-saturated solution showing cracks initiated in the surface of the LAZs.
Charpy impact values of the SA, H900, H1025 and LA specimens are 30, 10, 31 and 33 J, respectively. Peak-aged specimens had higher strength and lower toughness than the SA and overaged specimens. Impact energy of LA specimens is obviously higher than that of the H900 specimens. It also reveals that LA treatment can effectively improve the ability of 17-4 PH stainless steel in peak-aged condition to resist unstable brittle fracture.
3.4. Fatigue crack propagation
Fig. 7. Fatigue crack growth behavior for various specimens tested in laboratory air.
found to be more resistant to SCC than the peak-aged specimen under applied anodic potential [11]. Meanwhile, that the carbide precipitation would cause the depletion of Cr leading to a selective attack on the boundaries in the overaged specimen [11]. Thus, it was deduced that the amount of strain was higher within the relatively weak CR than the beside BM during tensile test, resulting in a premature fracture in the CR of the LA specimen. Examinations of the rectangular tensile fractured specimen tested in an H2S-saturated solution revealed that cracks were initiated in the LAZs around the specimen’s corners (Fig. 6), causing the
Fig. 7 shows the da/dN vs. DK curves of the various specimens tested in laboratory air. As shown in Fig. 2, the crack will propagate about 7.5 mm (involving 2 mm precrack) in the BM and then into the CR with 6 mm width. As the crack grows within the CR in an LA specimen, fatigue crack growth rates represent the overall effect of the LAZs and the sandwiched BM. When crack grew through the CR and into the BM again, fatigue crack growth rates of the LA specimen would be similar to those of the H900 specimen. The H900 specimen had the highest fatigue crack growth rates among the specimens, especially at low DKs. It was apparent that the retardation of crack growth in the region ahead of the CR, in which had the same microstructures and properties as those of the H900 specimen, was rather pronounced. Within the experimental DK range between 32.2MPa m and 44.5MPa m, it seemed that the H1025 specimen had slightly lower fatigue crack growth rates than the other specimens. In order to magnify the differences in fatigue crack growth rates in various regions of an LA specimen, constant DK tests were performed at a stress ratio of 0.1. Fig. 8 displays the results of constant DK tests with the crack propagation normal to LSD. Regardless of the testing DK range, the resistance to crack growth in the region in front of the CR was very obvious. At DK values of 26.8MPa m, similar fatigue crack growth rates between the CR and BM were obtained. At DK values of 39.2MPa m, the CR had slightly lower fatigue crack growth rates than the BM. Measurements of residual stresses by a hole-drilling strain gauge method conformed that principal tensile
70
L.W. Tsay et al. / Materials Science and Engineering A311 (2001) 64–73
residual stresses (+ 190MPa) were present in the centerline of the LAZs, and principal compressive residual stresses (− 98MPa) were generated in the BM 6 mm away from the centerline of the LAZs. The principal stress direction was parallel to the LSD. In a weld, the presence of welding residual stress is mainly due to the thermal shrinkage of the weld metal and heat-affected zone after welding. Shi et al. [18] proposed that the
Fig. 9. Fatigue crack growth behavior for various specimens tested in gaseous hydrogen.
Fig. 8. Crack growth rate (da/dN) versus crack length a tested under constant DK conditions, (a) DK= 26.8MPa m and (b) DK= 39.2MPa m.
partial relief and redistribution of residual stresses during crack growth in a weld caused the crack face to bend or rotate. As the crack grew transverse to the welding direction, residual compressive stresses would be induced ahead of the crack tip [18]. The influence of welding residual stresses on fatigue crack growth rates of A 514 steel welds has been investigated by Parry et al. [19]. The resistance to fatigue crack growth was higher for the as-welded weld metal and heat-affected zone than that of the BM. A considerable increase in fatigue crack growth rates was found after stress relief heat treatment. Residual tensile stresses are known to be detrimental to fatigue properties. The tough microstructure always had higher resistance to fatigue crack growth. It could be due to the beneficial effect of high toughness CR was offset partly by the deteriorated effect of residual tensile stresses. Hence, the CR had similar fatigue crack growth rates as the age-hardened BM under the same DK range. In contrast to this result, previous research work on the investigation of T-250 maraging steel annealed by laser [20] demonstrated that the retardation of crack growth within the CR of the LA specimen was very pronounced. Compressive residual stresses can enhance crack closure [21,22] and result in decreased fatigue crack growth rates. Therefore, lower fatigue crack growth rates in the region ahead of the CR in the LA specimen could be attributed to the presence of residual compressive stresses in this region. Fatigue crack growth behavior of CT specimens tested in gaseous hydrogen is shown in Fig. 9. It indicated that accelerated crack growth of H900-H,
L.W. Tsay et al. / Materials Science and Engineering A311 (2001) 64–73
H1025-H and LA-H specimens as compared with the H1025 specimen measured in air was found. The resistance to crack growth in the region ahead of the CR in the LA specimen was still obvious even when tested in gaseous hydrogen.
3.5. Fractographic obser6ations SEM fractographs of tensile fractured specimens tested in an H2S-saturated solution are shown in Fig. 10. Charpy impact and tensile fractured specimens tested in air displayed a ductile dimple fracture. Fracture appearance of tensile specimens tested in a corrosive environment revealed quasi-cleavage fracture in H900 and LA specimens (Fig. 10(a)). However, mixed mode fracture was observed for H1025 specimens (Fig. 10(b)). Microstructure observation of various specimens by transmission electron microscope had been performed but not shown here. For H1025 specimen, thin layer of reverted austenite precipitated in martensite lath boundaries, and pool-like reverted austenite together with o-Cu particles distributed homogeneously in the grain interior. Meanwhile, it was found that prior austenite grain boundaries delineated with relatively coarse o-Cu precipitates, as compared with those pre-
71
sented in the grain interior. It is known that grain boundaries are irreversible trapped sites for hydrogen. For the H1025 specimen tested in an H2S-saturated solution, the diffusion of hydrogen to grain boundaries, in which covered with dense o-Cu precipitates might reduce the cohesive strength of grain boundaries. Thus, the existence of grain boundary precipitates not only deteriorated tensile properties but also resulted in the formation of intergranular fracture of H1025 specimens. Microstructures of the alloy in SA condition revealed martensite laths had high density of dislocations. After aging at 482°C/1 h, very fine homogeneous precipitates were seen, as mentioned in Viswanathan investigation [5]. It was worth noting that few precipitates could be found at prior austenite grain boundaries in H900 specimens. Those fine and uniformly distributed precipitates provided a large amount of trapping sites and made the trap of hydrogen atoms more homogeneous. Thus, H900 specimens showed a higher resistance to SSCC than the other specimens. Decohesion between the precipitates and matrix interface, which was initiated by hydrogen embrittlement, led to the formation of quasi-cleavage fracture tested in an H2S solution. Fatigue fractographs of CT specimens tested in air and gaseous hydrogen are shown in Fig. 11. Fatiguefractured appearance of H900 and H1025 specimens indicated transgranular fatigue fracture under the experimental DK range (Fig. 11(a)). Although the decrease in fatigue crack growth rates in the region ahead of the CR in the LA specimen was very pronounced, fractographic examinations indicated no difference in fracture appearance as that found in the H900 specimen under the same DK values. Since the CR had similar fatigue crack growth rates as the H900 specimen, the LAZs also show transgranular fatigue fracture and clear fatigue striations at high DK values (Fig. 11(b)). A more marked change in fatigue fractography was observed for the specimens tested in gaseous hydrogen. Fracture surface consisted of large extent of quasicleavage fracture for all the specimens in gaseous hydrogen throughout the test (Fig. 11(c)). Accelerated crack growth in gaseous hydrogen for the various specimens was related to the brittle fracture features induced by hydrogen.
4. Conclusions
Fig. 10. SEM fractographs of tensile fractured specimens tested in an H2S-saturated solution showing (a) quasi-cleavage fracture in the H900 specimen and (b) mixed fracture modes in the H1025 specimen.
(1) Immersion test of various specimens in the H2Ssaturated solution indicated that SA, LA and H900 (peak-aged) specimens had the similar ability to resist hydrogen-induced cracking, meanwhile, surface cracks were initiated in the H1025 (overaged) specimen after a short time immersion in the H2S solution. H900 specimen strengthened by uniformly distributed fine precipi-
72
L.W. Tsay et al. / Materials Science and Engineering A311 (2001) 64–73
between the CR in a LA specimen and peak-aged specimens was not obvious. Such results could be due to the beneficial effect of high toughness CR was offset partly by the deteriorated effect of residual tensile stresses. (3) The H900 specimen had slightly higher fatigue crack growth rates than the other specimens, especially at low DK values. It was apparent that the retardation of crack growth in the region ahead of the CR in the LA specimen was very pronounced, regardless of testing environment. The presence of residual compressive stresses therein could account for the increased resistance to fatigue crack growth. (4) Fracture appearance of tensile fractured specimens tested in an H2S-saturated solution revealed quasi-cleavage fracture for H900 and LA specimens, and mixed mode fracture for the H1025 specimen. Although, the decrease in fatigue crack growth rates in the region ahead of the CR of the LA specimen was very obvious, fractographic examinations indicated no difference in fracture features from the H900 specimen under the same DK values. Large extent of quasi-cleavage fracture could account for enhanced crack growth of the various specimens tested in gaseous hydrogen.
Acknowledgements The authors gratefully acknowledge the support of the National Science Council of Republic of China (Contract no.87 NSC-2216-E-019-002).
References
Fig. 11. SEM fractographs of fatigue-fractured specimens showing (a) transgranular fatigue fracture of the H900 specimen in air, (b) transgranular fracture of the LAZs at high DK values in air, and (c) quasi-cleavage fracture of the H900 specimen in gaseous hydrogen.
tates had superior tensile properties than the other specimens. The increased hydrogen embrittlement susceptibility of H1025 specimens could be attributed to the presence of grain boundary precipitates. (2) Charpy impact energy of H900 specimens was lower than that of the SA and H1025 specimens. The results also indicated that laser-annealing treatment could effectively improve the ability of 17-4 PH stainless steel in peak-aged condition to resist unstable brittle fracture. Although the CR of the LA specimen showed much higher impact energy than the H900 specimen, the discrepancy in fatigue crack growth rates
[1] H.J. Rack, D. Kalish, Met. Trans. 5 (1974) 1595. [2] C.S. Carter, D.G. Farwick, A.M. Ross, J.M. Uchida, Corrosion 26 (1971) 190. [3] U. Kamachi Mudali, A.K. Bhaduri, J.B. Gnanamoorthy, Mats. Sci. Tech. 6 (1990) 475. [4] U.K. Viswanathan, P.K.K. Nayar, R. Krishnan, Mats. Sci. Tech. 5 (1989) 346. [5] U.K. Viswanathan, S. Banerjee, R. Krishnan, Mats. Sci. Eng. A104 (1988) 181. [6] R.M. Thompson, G.B. Kohut, D.R. Canfield, W.R. Bass, Corrosion 47 (1991) 216. [7] M.F. Ashby, K.E. Easterling, Acta Metall. 32 (1984) 1935. [8] H.J. Hegge, H. De Beurs, J. Noordhuis, J.Th.M. De Hosson, Met. Trans. 21A (1990) 987. [9] L.W. Tsay, Y.M. Li, C. Chen, S.W. Cheng, Int. J. Fatig. 14 (1992) 239. [10] V.M. Salinas-Bravo, J.G. Gonzalez-Rodriguez, Brit. Corr. J. 30 (1995) 77. [11] K.S. Raja, K. Prasad Rao, J. Mat. Sci. Lett. 12 (1993) 963. [12] B.J. Berkowitz, F.H. Heubaum, Corrosion 40 (1984) 240. [13] R. Gee, Z.Y. Ghen, Corr. Sci. 37 (1995) 2003. [14] A. Ikeda, T. Kaneko, Y. Ando, Corr. Sci. 27 (1987) 1099. [15] J.A. Beavers, G.H. Koch, Corrosion 48 (1992) 256. [16] S. Lin, Z.H.U. Ziyong, K.E. Wei, J. Mat. Sci. Tech. 12 (1996) 52.
L.W. Tsay et al. / Materials Science and Engineering A311 (2001) 64–73 [17] J.R. Scully, M.J. Cieslak, J.A. Van Den Avyle, Scri. Met. 31 (1994) 125. [18] Y.W. Shi, B.Y. Chen, J.X. Zhang, Eng. Fract. Mech. 36 (1990) 893. [19] M. Parry, H. Nordberg, R.W. Hertzberg, Weld. J. 51 (1972) 485s.
.
73
[20] L.W. Tsay, T.Y. Yang, Fat. Fract. Eng. Mat. Struc. 23 (2000) 325. [21] K.J. Kang, J.H. Song, Y.Y. Earmme, Fat. Farct. Eng. Mat. Struc. 13 (1990) 1. [22] M. Beghini, L. Bertini, E. Vitale, Fat. Fract. Eng. Mat. Struc. 17 (1994) 1443.
Embrittlement of laser surface-annealed 17-4 PH stainless steel L.W. Tsay *, T.Y. Yang, M.C. Young Institute of Materials Engineering, National Taiwan Ocean Uni6ersity, 2 Pei-Ning Road, Keelung 202, Taiwan, ROC Received 10 October 2000; received in revised form 18 December 2000
Abstract Stress corrosion cracking and fatigue crack growth behavior were determined in 17-4 PH stainless steel under various metallurgical conditions, including the H900 (482°C/1 h), H1025 (552°C/4 h) aging and laser surface annealing treatments. Peak-aged (H900) specimens locally irradiated by laser beam consisted of a portion of composite region (CR), in which comprised of soft laser-annealed (LA) zones on the outer surfaces and the hard base metal in between. Slow extension rate tensile tests were performed at room temperature in a saturated H2S solution to evaluate the hydrogen embrittlement susceptibility of various specimens. H900 specimens show an obvious improvement in impact toughness after irradiating by laser. Regardless of testing environments, H900 specimens exhibited better tensile properties than the other specimens. Experimental results also indicated that H900 specimens had the highest fatigue crack growth rates among the specimens, particularly at low stress intensity factor range. The retardation of crack growth in the region ahead of the CR in the LA specimens was rather pronounced. For compact tension specimens tested in gaseous hydrogen, enhanced crack growth was correlated with quasi-cleavage fracture in contrast to transgranular fatigue fracture in air. © 2001 Elsevier Science B.V. All rights reserved. Keywords: Fatigue crack growth; Hydrogen embrittlement; Composite region; Laser annealing treatment
1. Introduction 17-4 PH (AISI 630) steel is a precipitation-hardening martensitic stainless steel, which behaves high strength, high toughness and moderate corrosion resistance [1,2]. This steel can be used in differently aged conditions and is widely employed in various industries. The microstructure in the solution-annealed (SA) condition comprises a mixture of a martensite and 5 – 10 volume percent d ferrite, the latter elongated in the prior working direction [3]. A rapid increase in the hardness of the alloy during aging is owing to the initial formation of coherent copper-rich cluster in the peak-aged condition [4]. At temperature up to 600°C, overaging occurs resulting in the formation of incoherent o phase precipitates together with the transformation of a significant amount of martensite to austensite along martensite lath boundaries [5]. The microstructural changes not only alter the mechanical properties but also affect the corrosion resistance [3]. In general, the steel in the * Corresponding author. Fax: + 886-2-24625324. E-mail address: [email protected] (L.W. Tsay).
overaged condition has a higher resistance to sulfide stress corrosion cracking (SSCC) because of its lower hardness [6]. Both KIC and KISCC (in 3.5% aqueous sodium chloride solution) decreased inversely with strength level for high strength precipitation hardening stainless steels [2]. However, the precipitation of incoherent o phase in overaged specimens caused the depletion of Cu and Ni in the matrix, resulting in the enhanced electrochemical dissolution of the martensite matrix [3]. The use of laser irradiation in material processing has been increased in recent years. If the steel has a suitable carbon content after laser treatment, the near-surface region heated to above the austenizing temperature will be transformed to hard martensite [7]. The short thermal cycle of the laser processing can also result in the formation of softened zones in the overlapped areas after multiple laser passes [8]. The hardened welds could also be tempered effectively by controlling the power and irradiating area of the laser beam [9]. Depending on the thermal cycles at various locations, the fine precipitates in the surface region of the peak-aged specimen after laser irradiation may be dissolved into
0921-5093/01/$ - see front matter © 2001 Elsevier Science B.V. All rights reserved. PII: S 0 9 2 1 - 5 0 9 3 ( 0 1 ) 0 0 9 4 1 - 8
L.W. Tsay et al. / Materials Science and Engineering A311 (2001) 64–73
the matrix or coarsened. Hence, a specific depth of surface layer with lower hardness than the interior hard base metal (BM) is formed, and the soft surface layer is defined as the laser-annealed zone (LAZ). Locally symmetrical irradiation on the top and bottom surfaces of peak-aged specimens produced a composite region (CR), in which consisted of soft LAZs on the outer surfaces and untransformed BM sandwiched between them. In general, high strength steels are sensitive to hydrogen embrittlement and surface notches. It was reported [10] that both hydrogen embrittlement and anodic dissolution mechanisms are responsible for the stress corrosion cracking (SCC) of 17-4 PH steel in an acidic NaCl solution. Under applied anodic potentials, the 17-4 PH stainless steel in the SA state was found to be more resistant to SCC [11]. It is well known that an important factor affecting SSCC susceptibility is the strength level of the material, which is directly related to the material hardness. As defined in NACE standard MR0175-88, the maximum allowable hardness for 17-4 PH stainless steel is HRC33 to resist SSCC [6]. 17-4 PH stainless steel is extensively used in the power industry. In geothermal power application, structural components may suffer from hydrogen embrittlement, owing to the presence of H2S species in the environment. Laser irradiation on the peak-aged 17-4 PH stainless steel leads to the formation of soft surface layer in the SA condition with high ductility/toughness. The combination of soft LAZ on the surface and interior strong BM may not only possess adequate strength/ductility, but also reduce the susceptibility to SSCC and improved resistance to crack growth of the alloy. This study was undertaken with the objective of characterizing the influence of aging condition and hydrogen charging on the tensile and fatigue crack growth behavior in 17-4 PH stainless steel aged at various conditions. It is reported that SSCC was a manifestation of hydrogen embrittlement [12,13]. Slow extension rate tensile tests were usually performed to evaluate the susceptibility to hydrogen embrittlement of various specimens in an H2S-saturated NACE solution [14,15]. Impact energy of the specimens measured at room temperature was used to assess the influence of laser irradiation on the improved resistance to unstable brittle fracture, and the results were compared with aged specimens. It is found that few research works have been performed to study the fatigue crack growth behavior in 17-4 PH stainless steel aged at various conditions. Compact-tension (CT) specimens tested in air and gaseous hydrogen were used to determine the fatigue crack growth rates of various specimens, and the effect of hydrogen embrittlement on accelerating crack growth was evaluated. Two types of fatigue tests, constant load and constant DK were conducted in this
65
work. Fractographic observations were performed to correlate the microstructures and changes in properties with fracture features.
2. Materials and experimental procedures The chemical composition in wt.% of the 17-4 PH stainless steel used in this investigation in the plate form of 5.6 mm thick was 15.71Cr, 4.35Ni, 3.46Cu, 0.04C, 0.7Si, 0.38Mn, 0.025P, 0.013S and balance Fe. The round bar of 15 mm diameter was composed of (wt.%) 16.60Cr, 4.26Ni, 3.50Cu, 0.046C, 0.52Si, 0.6Mn, 0.03P, 0.043S and balance Fe. The alloy was SA at 1038°C for 1 h and then air-cooled to room temperature. The aging treatments were performed either at 482°C for 1 h (H900) or 552°C for 4 h (H1025). Some of the H900 specimens were subjected to a laser-annealed (LA) treatment. In this investigation, laser beam integrators (6× 6 mm2) and rectangular beams (6×25 mm2) were used to provide uniform distribution of energy over a relatively large area, resulting in the formation of a more homogeneously annealed region without surface-melting of the specimen. A Rofin –Sinar 5 KW CO2 laser connected with a CNC table was utilized for laser treatment. Steel plates aged at 482°C were processed by laser at 2000 W and 800 mm min − 1 scan rate and the round specimens were laser-treated at 600 W with 2 turns s − 1 for 5 s. The hardness of the steel plate in various conditions, including SA, as well as H900 and H1025 aging treatment, was measured by Rockwell hardness tester (HRC). After laser treatment, some work pieces were sectioned transverse to the laser scan direction (LSD), then polished and etched in a Vilella’s reagent. Microhardness measurements across the LAZ was taken via a Vickers microhardness tester. Fig. 1 shows the schematic dimensions of tensile and Charpy impact specimens employed in this study. Two types of tensile specimens with a gauge length of 25 mm are shown in Fig. 1. The tensile properties of various specimens tested in air were determined at an ordinary strain rate of 5×10 − 3 s − 1. In order to evaluate the hydrogen embrittlement susceptibility of various specimens, slow extension rate tensile tests were carried out at room temperature in an H2S-saturated solution with a constant strain rate of 5× 10 − 6 s − 1. The test solution prepared according to the NACE standard (TM-01/7786) was purged with N2 for at least 1 h to get rid off O2 before testing and then H2S was bubbled continuously to ensure that the solution was saturated with H2S. The results represent the average of at least three specimens for each testing condition. Charpy impact specimens (Fig. 1(c)) are made according to the specifications of ASTM E23 (subsize, 5 mm thick) fractured at room temperature.
66
L.W. Tsay et al. / Materials Science and Engineering A311 (2001) 64–73
The fatigue crack growth tests were performed at room temperature in a computerized hydraulic testing machine (MTS 810). CT specimens are used and the experimental procedures meet the standard of ASTM E647-91 specification. The geometry of the specimens used in this study is shown in Fig. 2, in which the crack growth direction is normal to the LSD. The testing software (759.40 crack growth test) developed by MTS can offer a variety of testing requirements including constant load, constant DK and K-controlled step testing conditions. In order to magnify the differences in fatigue crack growth rates in various regions of a LA specimen, constant DK tests were performed at a stress ratio of 0.1. In addition to the CT specimen that were tested in air, the effect of hydrogen embrittlement on
Fig. 2. Schematic diagram showing the dimensions of CT specimen.
accelerating the crack growth was investigated by testing the CT specimens inside a stainless steel chamber with 2 atm. hydrogen gas. In such a condition, a capital ‘‘H’’ is added to the last symbol of the designated specimen, e.g. H900-H represents the hydrogen-charged H900 specimen. Fractographic observations of the fractured specimens were examined by a Hitachi 4100 scanning electron microscope (SEM), with emphasis paid to the change in fracture modes. Residual stresses measurement (ASTM E837-92) was performed by a modified hole-drilling strain gauge method. Three-element strain gauge rosettes (Measurements Group, CEA-06-062UL-120) were applied onto the CT specimen at the centerline of the LAZ and 6 mm away from the centerline (in the BM). A hole of 1.6 mm diameter was drilled by an air turbine hole-drilled machine in the center of the rosette. Relaxed strains in terms of residual stresses were measured by a strain indicator.
3. Results and discussion
3.1. Hardness measurements and metallographic obser6ations
Fig. 1. Schematic diagrams showing the dimensions of (a) round tensile specimen, (b) rectangular tensile specimen and (c) Charpy impact specimen; dimension in mm.
Both the Rockwell and Vickers hardness testers were used to measure the hardness of 17-4 PH steel plates aged in various conditions. The initial hardness is of about HRC32.1 (HV370) for the specimen in the SA condition. After aging at 482°C for 1 h (H900), the precipitation reaction resulted in an increase in hardness to HRC44.5 (HV465). Aging at a higher tempera-
L.W. Tsay et al. / Materials Science and Engineering A311 (2001) 64–73
ture, e.g., 552°C (H1025), caused the hardness to decline to HRC37.2 (HV389). Fig. 3 displays schematic diagram of microhardness distribution in LA specimens measured at the center of the LA region for the round and rectangular tensile specimens. A decrease in surface hardness to HV370 similar to the hardness of the SA specimen was noted. It suggested that the near-surface
67
Table 1 Tensile properties of 17-4 PH steel in various conditions tested in air and H2S solution for round tensile specimensb Materials Properties/Condition
UTS (MPa) Elongation (%) Reduction in area (%) Fracture location
H900
H900+LA
In Aira
In H2Sb
In Air
In H2S
1273 19.3 30
1157 7.5 15.2
1181 10.3 27.1
1099 6.5 15.0
–
–
CR
CR
a UTS – ultimate tensile strength, CR – composite region. In Air – tested in laboratory air with the strain rate of 5×10−3 s−1. b In H2S – tested in an H2S-saturated solution with the strain rate of 5×10−6 s−1.
Fig. 3. Schematic diagram showing microhardness distribution in the rectangular and round LA specimens.
region after laser irradiation should be transformed back to the SA condition. In addition, the variation in microhardness distribution in the LA specimen could be attributed to the complex matrix/precipitate reactions resulting in a continuous phase transformation within this narrow heat-affected region. Fig. 4 reveals the metallographs of the various specimens. The microstructures of the H900 specimen consisted of a mixture of martensite and elongated d ferrite oriented along the rolling direction (Fig. 4(a)). The SEM/EDX spectrum revealed that the Cr concentration in the d ferrite was slightly higher than that in the matrix. It has been found that the near-surface region in the LAZ contained less stringer d ferrite (upper portion of Fig. 4(b)). It should be due to the high temperature thermal cycle associated with laser treatment resulting in the dissolution of d ferrite into the matrix. Hence, elongated d ferrite was refined at the near-surface region. It had been pointed out previously [3] that the compositional difference in Cr could account for the localized attack at the d ferrite/matrix interface. Thus, improvement in corrosion resistance of the LAZ could be anticipated, owing to the dissolution of d ferrite as well as the reduction of d ferrite/matrix interface.
3.2. Tensile properties
Fig. 4. Metallographs showing (a) the elongated d ferrite in the steel plate and (b) refined d ferrite in the near-surface region of an LA specimen. Arrows indicating the stringer d ferrite.
Table 1 and Table 2 present the tensile properties of 17-4 PH stainless steel aged in various conditions tested in air and H2S solution for round and rectangular tensile specimens, respectively. In case of the round tensile specimen annealed by laser, a specific depth of soft surface layer (i.e. LAZ) enclosing the interior hard BM is shown in Fig. 1(a). The CR of the round tensile specimen included the LAZ at the periphery and interior BM. As shown in Fig. 1(b), semi-elliptical LAZs were located on the top and bottom surfaces and BM sandwiched between them in rectangular tensile speci-
68
L.W. Tsay et al. / Materials Science and Engineering A311 (2001) 64–73
Table 2 Tensile properties of 17-4 PH steel in various conditions tested in air and H2S solution for rectangular tensile specimens Materials Properties/Condition
UTS (MPa) Elongation (%) Reduction in area (%) Fracture location
H900
H1025
H900+LA
In Air
In H2S
In Air
In H2S
In Air
In H2S
1356 19.2 50.1 –
948 3.6 2.2 –
1087 21.7 57 –
822 4.2 3.6 –
1234 14.3 49.2 CR
822 4.1 3.0 CR
mens. In this case, the BM in the CR was not completely enclosed by the LAZs. It was noted that a slightly different degree of embrittlement was found for those two types of CR testing in a corrosive environment. The sandwiched BM in the CR of a rectangular tensile specimen would directly immerse into H2S solution, in contrast to that of round tensile specimens, the BM was surrounded by the LAZ. For the round tensile specimen tested in air, the H900 specimen had slightly better tensile properties than the LA specimen. When tested in an H2S-saturated solution, both the H900 and LA specimen showed a slight decrease in strength but greater reduction in ductility as compared to that tested in air. The fracture location of the LA specimen, regardless of testing environment, was in the CR. As given in Table 2, rectangular tensile specimens show the same change of trend in tensile properties as round tensile specimens. H900 specimens possessed the best tensile properties among specimens tested in air, and H1025 specimens had lower strength than the other specimens. When tested in an H2S-saturated solution, H900 specimens still exhibited the highest strength among the various specimens. Meanwhile, the results also indicated that specimen’s geometry had a great influence on the tensile strength of the specimens tested in an H2S-saturated solution. It was known that the fracture stress of the material in hydrogen decreased as the notch severity was increased [16]. A more obvious degradation in tensile strength for a rectangular specimen could be attributed to the more inhomogeneous distribution of tensile stress over the cross-section of the specimen. Subsequently, all the specimens were immersed in the saturated H2S solution and the surface cracks were inspected after a period of time. The optical micrographs revealing the surface morphology of immersed specimens are shown in Fig. 5. It was found that surface cracks were initiated in the H1025 specimen after being immersed in an H2S-saturated solution for 7 h (Fig. 5(a)) but no detectable surface cracks were found in other specimens. Charging the specimens in H2S for 20 h yielded the same surface defects on the specimens (LA, SA and H900 specimens) i.e., corrosion
pits were formed and eventually coalesced into cracks (Fig. 5(b)). Such results demonstrated that SA, LA and H900 specimens had the similar ability to resist hydrogen-induced cracking, meanwhile, H1025 specimen was more sensitive to hydrogen embrittlement. SSCC susceptibility in various region of a weld has been found to vary with the applied stress level [14]. It was reported [14] that SSCC could occur at the relatively soft region under a high stress level. In this work, the immersion tests showed that the H900 specimen revealed similar susceptibility to hydrogen-induced cracking as the LAZs within the testing time interval. It was pointed out previously that SA specimen was
Fig. 5. Optical micrographs revealing (a) surface cracks in the H1025 specimen immersed in an H2S-saturated solution for 7 h, and (b) corrosion pits in the H900 specimen immersed in an H2S-saturated solution for 20 h.
L.W. Tsay et al. / Materials Science and Engineering A311 (2001) 64–73
69
decrease in tensile strength and ductility of LA specimens. It has also been reported that d ferrite in precipitation age-hardened stainless steels will increase the susceptibility to hydrogen embrittlement [17]. When the testing strain rate was decreased or the severe notch was introduced, the CR might show lower susceptibility to hydrogen embrittlement and better resistance to SSCC than the hard BM, which was related to the less stringer d ferrite and lower surface hardness.
3.3. Charpy impact test
Fig. 6. Macroscopic photograph of tensile fractured specimens tested in an H2S-saturated solution showing cracks initiated in the surface of the LAZs.
Charpy impact values of the SA, H900, H1025 and LA specimens are 30, 10, 31 and 33 J, respectively. Peak-aged specimens had higher strength and lower toughness than the SA and overaged specimens. Impact energy of LA specimens is obviously higher than that of the H900 specimens. It also reveals that LA treatment can effectively improve the ability of 17-4 PH stainless steel in peak-aged condition to resist unstable brittle fracture.
3.4. Fatigue crack propagation
Fig. 7. Fatigue crack growth behavior for various specimens tested in laboratory air.
found to be more resistant to SCC than the peak-aged specimen under applied anodic potential [11]. Meanwhile, that the carbide precipitation would cause the depletion of Cr leading to a selective attack on the boundaries in the overaged specimen [11]. Thus, it was deduced that the amount of strain was higher within the relatively weak CR than the beside BM during tensile test, resulting in a premature fracture in the CR of the LA specimen. Examinations of the rectangular tensile fractured specimen tested in an H2S-saturated solution revealed that cracks were initiated in the LAZs around the specimen’s corners (Fig. 6), causing the
Fig. 7 shows the da/dN vs. DK curves of the various specimens tested in laboratory air. As shown in Fig. 2, the crack will propagate about 7.5 mm (involving 2 mm precrack) in the BM and then into the CR with 6 mm width. As the crack grows within the CR in an LA specimen, fatigue crack growth rates represent the overall effect of the LAZs and the sandwiched BM. When crack grew through the CR and into the BM again, fatigue crack growth rates of the LA specimen would be similar to those of the H900 specimen. The H900 specimen had the highest fatigue crack growth rates among the specimens, especially at low DKs. It was apparent that the retardation of crack growth in the region ahead of the CR, in which had the same microstructures and properties as those of the H900 specimen, was rather pronounced. Within the experimental DK range between 32.2MPa m and 44.5MPa m, it seemed that the H1025 specimen had slightly lower fatigue crack growth rates than the other specimens. In order to magnify the differences in fatigue crack growth rates in various regions of an LA specimen, constant DK tests were performed at a stress ratio of 0.1. Fig. 8 displays the results of constant DK tests with the crack propagation normal to LSD. Regardless of the testing DK range, the resistance to crack growth in the region in front of the CR was very obvious. At DK values of 26.8MPa m, similar fatigue crack growth rates between the CR and BM were obtained. At DK values of 39.2MPa m, the CR had slightly lower fatigue crack growth rates than the BM. Measurements of residual stresses by a hole-drilling strain gauge method conformed that principal tensile
70
L.W. Tsay et al. / Materials Science and Engineering A311 (2001) 64–73
residual stresses (+ 190MPa) were present in the centerline of the LAZs, and principal compressive residual stresses (− 98MPa) were generated in the BM 6 mm away from the centerline of the LAZs. The principal stress direction was parallel to the LSD. In a weld, the presence of welding residual stress is mainly due to the thermal shrinkage of the weld metal and heat-affected zone after welding. Shi et al. [18] proposed that the
Fig. 9. Fatigue crack growth behavior for various specimens tested in gaseous hydrogen.
Fig. 8. Crack growth rate (da/dN) versus crack length a tested under constant DK conditions, (a) DK= 26.8MPa m and (b) DK= 39.2MPa m.
partial relief and redistribution of residual stresses during crack growth in a weld caused the crack face to bend or rotate. As the crack grew transverse to the welding direction, residual compressive stresses would be induced ahead of the crack tip [18]. The influence of welding residual stresses on fatigue crack growth rates of A 514 steel welds has been investigated by Parry et al. [19]. The resistance to fatigue crack growth was higher for the as-welded weld metal and heat-affected zone than that of the BM. A considerable increase in fatigue crack growth rates was found after stress relief heat treatment. Residual tensile stresses are known to be detrimental to fatigue properties. The tough microstructure always had higher resistance to fatigue crack growth. It could be due to the beneficial effect of high toughness CR was offset partly by the deteriorated effect of residual tensile stresses. Hence, the CR had similar fatigue crack growth rates as the age-hardened BM under the same DK range. In contrast to this result, previous research work on the investigation of T-250 maraging steel annealed by laser [20] demonstrated that the retardation of crack growth within the CR of the LA specimen was very pronounced. Compressive residual stresses can enhance crack closure [21,22] and result in decreased fatigue crack growth rates. Therefore, lower fatigue crack growth rates in the region ahead of the CR in the LA specimen could be attributed to the presence of residual compressive stresses in this region. Fatigue crack growth behavior of CT specimens tested in gaseous hydrogen is shown in Fig. 9. It indicated that accelerated crack growth of H900-H,
L.W. Tsay et al. / Materials Science and Engineering A311 (2001) 64–73
H1025-H and LA-H specimens as compared with the H1025 specimen measured in air was found. The resistance to crack growth in the region ahead of the CR in the LA specimen was still obvious even when tested in gaseous hydrogen.
3.5. Fractographic obser6ations SEM fractographs of tensile fractured specimens tested in an H2S-saturated solution are shown in Fig. 10. Charpy impact and tensile fractured specimens tested in air displayed a ductile dimple fracture. Fracture appearance of tensile specimens tested in a corrosive environment revealed quasi-cleavage fracture in H900 and LA specimens (Fig. 10(a)). However, mixed mode fracture was observed for H1025 specimens (Fig. 10(b)). Microstructure observation of various specimens by transmission electron microscope had been performed but not shown here. For H1025 specimen, thin layer of reverted austenite precipitated in martensite lath boundaries, and pool-like reverted austenite together with o-Cu particles distributed homogeneously in the grain interior. Meanwhile, it was found that prior austenite grain boundaries delineated with relatively coarse o-Cu precipitates, as compared with those pre-
71
sented in the grain interior. It is known that grain boundaries are irreversible trapped sites for hydrogen. For the H1025 specimen tested in an H2S-saturated solution, the diffusion of hydrogen to grain boundaries, in which covered with dense o-Cu precipitates might reduce the cohesive strength of grain boundaries. Thus, the existence of grain boundary precipitates not only deteriorated tensile properties but also resulted in the formation of intergranular fracture of H1025 specimens. Microstructures of the alloy in SA condition revealed martensite laths had high density of dislocations. After aging at 482°C/1 h, very fine homogeneous precipitates were seen, as mentioned in Viswanathan investigation [5]. It was worth noting that few precipitates could be found at prior austenite grain boundaries in H900 specimens. Those fine and uniformly distributed precipitates provided a large amount of trapping sites and made the trap of hydrogen atoms more homogeneous. Thus, H900 specimens showed a higher resistance to SSCC than the other specimens. Decohesion between the precipitates and matrix interface, which was initiated by hydrogen embrittlement, led to the formation of quasi-cleavage fracture tested in an H2S solution. Fatigue fractographs of CT specimens tested in air and gaseous hydrogen are shown in Fig. 11. Fatiguefractured appearance of H900 and H1025 specimens indicated transgranular fatigue fracture under the experimental DK range (Fig. 11(a)). Although the decrease in fatigue crack growth rates in the region ahead of the CR in the LA specimen was very pronounced, fractographic examinations indicated no difference in fracture appearance as that found in the H900 specimen under the same DK values. Since the CR had similar fatigue crack growth rates as the H900 specimen, the LAZs also show transgranular fatigue fracture and clear fatigue striations at high DK values (Fig. 11(b)). A more marked change in fatigue fractography was observed for the specimens tested in gaseous hydrogen. Fracture surface consisted of large extent of quasicleavage fracture for all the specimens in gaseous hydrogen throughout the test (Fig. 11(c)). Accelerated crack growth in gaseous hydrogen for the various specimens was related to the brittle fracture features induced by hydrogen.
4. Conclusions
Fig. 10. SEM fractographs of tensile fractured specimens tested in an H2S-saturated solution showing (a) quasi-cleavage fracture in the H900 specimen and (b) mixed fracture modes in the H1025 specimen.
(1) Immersion test of various specimens in the H2Ssaturated solution indicated that SA, LA and H900 (peak-aged) specimens had the similar ability to resist hydrogen-induced cracking, meanwhile, surface cracks were initiated in the H1025 (overaged) specimen after a short time immersion in the H2S solution. H900 specimen strengthened by uniformly distributed fine precipi-
72
L.W. Tsay et al. / Materials Science and Engineering A311 (2001) 64–73
between the CR in a LA specimen and peak-aged specimens was not obvious. Such results could be due to the beneficial effect of high toughness CR was offset partly by the deteriorated effect of residual tensile stresses. (3) The H900 specimen had slightly higher fatigue crack growth rates than the other specimens, especially at low DK values. It was apparent that the retardation of crack growth in the region ahead of the CR in the LA specimen was very pronounced, regardless of testing environment. The presence of residual compressive stresses therein could account for the increased resistance to fatigue crack growth. (4) Fracture appearance of tensile fractured specimens tested in an H2S-saturated solution revealed quasi-cleavage fracture for H900 and LA specimens, and mixed mode fracture for the H1025 specimen. Although, the decrease in fatigue crack growth rates in the region ahead of the CR of the LA specimen was very obvious, fractographic examinations indicated no difference in fracture features from the H900 specimen under the same DK values. Large extent of quasi-cleavage fracture could account for enhanced crack growth of the various specimens tested in gaseous hydrogen.
Acknowledgements The authors gratefully acknowledge the support of the National Science Council of Republic of China (Contract no.87 NSC-2216-E-019-002).
References
Fig. 11. SEM fractographs of fatigue-fractured specimens showing (a) transgranular fatigue fracture of the H900 specimen in air, (b) transgranular fracture of the LAZs at high DK values in air, and (c) quasi-cleavage fracture of the H900 specimen in gaseous hydrogen.
tates had superior tensile properties than the other specimens. The increased hydrogen embrittlement susceptibility of H1025 specimens could be attributed to the presence of grain boundary precipitates. (2) Charpy impact energy of H900 specimens was lower than that of the SA and H1025 specimens. The results also indicated that laser-annealing treatment could effectively improve the ability of 17-4 PH stainless steel in peak-aged condition to resist unstable brittle fracture. Although the CR of the LA specimen showed much higher impact energy than the H900 specimen, the discrepancy in fatigue crack growth rates
[1] H.J. Rack, D. Kalish, Met. Trans. 5 (1974) 1595. [2] C.S. Carter, D.G. Farwick, A.M. Ross, J.M. Uchida, Corrosion 26 (1971) 190. [3] U. Kamachi Mudali, A.K. Bhaduri, J.B. Gnanamoorthy, Mats. Sci. Tech. 6 (1990) 475. [4] U.K. Viswanathan, P.K.K. Nayar, R. Krishnan, Mats. Sci. Tech. 5 (1989) 346. [5] U.K. Viswanathan, S. Banerjee, R. Krishnan, Mats. Sci. Eng. A104 (1988) 181. [6] R.M. Thompson, G.B. Kohut, D.R. Canfield, W.R. Bass, Corrosion 47 (1991) 216. [7] M.F. Ashby, K.E. Easterling, Acta Metall. 32 (1984) 1935. [8] H.J. Hegge, H. De Beurs, J. Noordhuis, J.Th.M. De Hosson, Met. Trans. 21A (1990) 987. [9] L.W. Tsay, Y.M. Li, C. Chen, S.W. Cheng, Int. J. Fatig. 14 (1992) 239. [10] V.M. Salinas-Bravo, J.G. Gonzalez-Rodriguez, Brit. Corr. J. 30 (1995) 77. [11] K.S. Raja, K. Prasad Rao, J. Mat. Sci. Lett. 12 (1993) 963. [12] B.J. Berkowitz, F.H. Heubaum, Corrosion 40 (1984) 240. [13] R. Gee, Z.Y. Ghen, Corr. Sci. 37 (1995) 2003. [14] A. Ikeda, T. Kaneko, Y. Ando, Corr. Sci. 27 (1987) 1099. [15] J.A. Beavers, G.H. Koch, Corrosion 48 (1992) 256. [16] S. Lin, Z.H.U. Ziyong, K.E. Wei, J. Mat. Sci. Tech. 12 (1996) 52.
L.W. Tsay et al. / Materials Science and Engineering A311 (2001) 64–73 [17] J.R. Scully, M.J. Cieslak, J.A. Van Den Avyle, Scri. Met. 31 (1994) 125. [18] Y.W. Shi, B.Y. Chen, J.X. Zhang, Eng. Fract. Mech. 36 (1990) 893. [19] M. Parry, H. Nordberg, R.W. Hertzberg, Weld. J. 51 (1972) 485s.
.
73
[20] L.W. Tsay, T.Y. Yang, Fat. Fract. Eng. Mat. Struc. 23 (2000) 325. [21] K.J. Kang, J.H. Song, Y.Y. Earmme, Fat. Farct. Eng. Mat. Struc. 13 (1990) 1. [22] M. Beghini, L. Bertini, E. Vitale, Fat. Fract. Eng. Mat. Struc. 17 (1994) 1443.