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Notch tensile properties of laser-surface-annealed 17-4 PH stainless steel in hydrogen-related environments
Corrosion Science 44 (2002) 2101–2118 www.elsevier.com/locate/corsci
Notch tensile properties of laser-surface-annealed 17-4 PH stainless steel in hydrogen-related environments L.W. Tsay a
a,*
, W.C. Lee a, R.K. Shiue b, J.K. Wu
a
Institute of Materials Engineering, National Taiwan Ocean University, Pei-Ning Road, Keelung 202, Taiwan, ROC b Department of Materials Science and Engineering, National Dong Hwa University, Hualien 974, Taiwan, ROC Received 21 September 2001; accepted 27 December 2001
Abstract Slow displacement rate tensile tests were performed to determine the notched tensile strength (NTS) of 17-4 PH stainless steel with various microstructures in hydrogen-related environments. Solution-annealed (SA), peak-aged (H900), over-aged (H1025), and laserannealed (LA) specimens were included in the study. Based on the results of NTS in air, the NTS loss in both gaseous hydrogen and H2 S-saturated solution was used to access the detrimental effects of hydrogen in 17-4 PH steel subjected to different treatments. Electrochemical permeation tests were also employed to determine the hydrogen permeation characteristics of the 17-4 PH steel plate with various microstructures. The result indicates that all the specimens have low NTS loss in gaseous hydrogen but significantly suffer from sulfide stress corrosion cracking (SCC), especially for the soft SA specimen. It was deduced that high hydrogen diffusivity and less trapped hydrogen atoms in the SA matrix provided rapid transport of massive hydrogen atoms into highly stressed region, and deteriorated the NTS tested in the saturated sulfide solution. On the other hand, H1025 specimen consists of the blocky austenite together with Cu-rich precipitates uniformly distributed in the grain interior; dense and coarse precipitates are also observed along prior austenite grain boundaries. Hydrogen atoms tend to be trapped along grain boundaries, and lead to the formation of intergranular fracture for H1025 specimen tested in the H2 S solution. Fine and homogeneously distributed precipitates in the H900 matrix result in uniformly trapping of hydrogen atoms, so it behaves superior properties than other specimens. The decohesion of precipitate/matrix interfaces induces quasi-cleavage fracture of the H900 specimen tested in H2 S solution. Finally, the application
*
Corresponding author. Tel.: +886-2-24622192; fax: +886-2-24625324. E-mail addresses: [email protected] (L.W. Tsay), [email protected] (R.K. Shiue).
0010-938X/02/$ - see front matter Ó 2002 Elsevier Science Ltd. All rights reserved. PII: S 0 0 1 0 - 9 3 8 X ( 0 2 ) 0 0 0 2 3 - 9
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of laser-annealing treatment on the H900 specimen cannot improve its resistance to sulfide SCC, because the laser-annealed zone is susceptible to hydrogen embrittlement in the H2 S solution. Ó 2002 Elsevier Science Ltd. All rights reserved. Keywords: C. Hydrogen embrittlement; Sulfide stress corrosion cracking; 17-4 PH stainless steels; Notch tensile strength
1. Introduction The 17-4 PH (AISI 630) steel is a precipitation-hardened martensitic stainless steel with high strength as well as moderate corrosion resistance, and it has been widely used in various industries [1,2]. After solution-annealed (SA) treatment, 17-4 PH stainless steel can be hardened by various aging treatments. The formation of coherent copper-rich clusters in the peak-aged condition is primarily responsible for the rapid increase of its hardness [3]. As the aging temperature is further increased above 600 °C, both the formation of incoherent e precipitates in the matrix and a significant amount of martensite transform into austenite along martensite lath boundaries [3]. The precipitation of incoherent e phase during over-aging will have depleted substantial amounts of Cu and Ni from the matrix, resulting in enhanced electrochemical dissolution of the martensite matrix [4]. It is well known that surface properties of many structural alloys can be greatly improved by the laser surface treatment [5,6]. After irradiating by the laser beam, the age-hardened alloy forms a soft ductile surface region in SA condition, which is socalled laser-annealed zone (LAZ) [7]. The composite region is characterized by such soft LAZs on its outer surfaces and the untransformed base metal sandwiched between two LAZs [7]. Meanwhile, it is reported that the improvement of fatigue crack growth is very pronounced in the region preceding the composite region for peakaged 17-4 PH stainless steel subjected to laser-surface-annealed treatment [8]. The stress corrosion cracking (SCC) of an alloy is strongly related to three important factors, including the microstructure of the alloy, tri-axial stress state of the specimen and environment of the alloy therein. In general, the hardened low alloy steel applied in hydrogen-related environments under high tensile stress is much more susceptible to SCC than the highly tempered one employed in the same environment under compressive stress. Traditionally, the hydrogen embrittlement susceptibility of high strength alloys can be greatly improved by decreasing the yield strength of the alloy. The age-hardened alloy in over-aged condition is usually more resistant to hydrogen embrittlement than that in peak-aged condition. 17-4 PH stainless steel is extensively used in the power industry. In geothermal power applications, the presence of H2 S in the environment may result in severe hydrogen embrittlement of the alloy. It is reported that both hydrogen embrittlement and anodic dissolution mechanism are responsible for SCC of the 17-4 PH stainless steel in an acidic NaCl solution [9]. The 17-4 PH stainless steel in the SA state is found to be very resistant to SCC in anodic potential environment [10]. Thus, it was deduced that the peak-aged 17-4 PH stainless steel with LAZs on the surfaces might
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not only possess adequate strength but also exhibited improved resistance to hydrogen embrittlement. Fractographs of many low alloy steels susceptible to hydrogen embrittlement depends on its microstructure, crack-tip stress intensity and hydrogen concentration around the crack tip [11]. Moreover, the susceptible alloy suffered from hydrogen embrittlement usually associates with the change of fracture modes [12,13]. SCC of low strength steels is characterized by transgranular fracture in contrast to intergranular fracture of high strength ones [12–14]. It has also been pointed out that hydrogen atoms are prone to diffuse into locations with high hydrostatic stress [15– 17]. In addition, hydrogen atoms tend to be accumulated and embrittle the elastic– plastic boundary of the high strength steel [17]. Consequently, notched tensile test under a slow displacement rate is a suitable experimental process in evaluating the detrimental effect of hydrogen to the alloy with various microstructures. The purpose of this investigation was focused on the influence of hydrogen embrittlement on 17-4 PH steel with various microstructures. Both degradation of notched tensile strength (NTS) and related fracture characteristics of 17-4 PH stainless steel were studied extensively. Slow displacement rate tensile tests were performed to enhance the interaction between the alloy and the hydrogen-related environment. The detrimental effect of hydrogen was accessed by comparing the NTS in air, gaseous hydrogen and H2 S-saturated NACE solution, respectively. Electrochemical permeation was employed to determine hydrogen permeation properties of the specimens aged at different conditions. Fractographic observations were carried out on the tensile fractured specimen focused on the neighborhood of the crack initiation sites, in which the changes in fracture morphology were usually observed. The deterioration of NTS was strongly related to the extent of embrittled area on its fracture surface. The influence of hydrogen embrittlement on the NTS of the specimens was explained on the basis of hydrogen/microstructure interactions, i.e. in terms of hydrogen diffusivity and hydrogen solubility of the alloy.
2. Experimental procedures The chemical composition in wt.% of the 17-4 PH stainless steel was 15.71Cr, 4.35Ni, 3.46Cu, 0.70Si, 0.04C, 0.38Mn, 0.025P, 0.013S and balance Fe. The asreceived steel plate with 5.6 mm in thickness was ground into 5.0 mm thick. It was firstly austenized at 1038 °C for 1 h and subsequently cooled by purging highpressure nitrogen; it was specified as the SA specimen with the hardness of HRC 32. The following aging treatment was performed at 482 °C for 1 h (H900) and 552 °C for 4 h (H1025), respectively. The hardness of peak-aged specimen (H900) was HRC 42, and the hardness of over-aged specimen (H1025) was HRC 34. Laser surface treatments were carried out by using a Rofin-Sinar 5 kW CO2 laser. The size of rectangular laser beam on the specimen was about 6 mm 25 mm with uniform energy intensity within the irradiated area. In order to enhance the absorptivity of laser irradiated surface, the specimen was sprayed by carbon black paint prior to laser treatment. Some H900 specimens were subjected to LA treatment.
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Fig. 1. Schematic diagram showing the dimensions of notched tensile specimen used in the experiment.
After laser irradiation, the specimen comprised of the composite region was recognized as LA specimens. The laser energy used in the experiment was 2000 W, and the scan rate of laser beam was set at 600 mm/min throughout the experiment. The microhardness measurement was performed along the centerline of the LA specimen. The schematic dimensions of notched tensile specimen used in this research are shown in Fig. 1. It is noted that the central shaded area in the figure is the LAZ. The notch tip radius shown in Fig. 1 is about 100 lm. Initial tensile tests were performed in air with the controlled crosshead speed of 0.75 mm/min. The effect of hydrogen embrittlement on the deterioration of NTS was evaluated by installing the specimen into a Teflon chamber with controlled atmosphere, i.e., gaseous hydrogen at 2 atm or saturated H2 S solution. It was reported that sulfide SCC was a manifestation of hydrogen embrittlement [18,19]. The H2 S solution was prepared according to NACE standard (TM-01-77-86). The solution was purged with nitrogen gas for at least 1 h to remove O2 before testing, and H2 S was continuously bubbled into the solution to ensure the solution properly saturated. In order to clarify the effect of hydrogen embrittlement on the alloy, a much slower crosshead speed of 7:5 103 mm/min was chosen during the tensile test. All NTS data were averaged from at least three specimens for each test condition. The degree of susceptibility to hydrogen embrittlement of various specimens can be expressed in the percentage loss of NTS under hydrogen charging as given below: NTS loss ð%Þ ¼
NTS ðin airÞ NTS ðin H2 or H2 SÞ 100% NTS ðin airÞ
ð1Þ
Tensile fractured appearance of various specimens was examined by a Hitachi S4100 field emission scanning electron microscope (SEM). Thin foil specimens for transmission electron microscope (TEM) observation were examined by a JEOL2000EX microscope with the operation voltage of 200 kV. Electrochemical permeation technique originally developed by Devanathan and Stachurski was employed to determine the hydrogen permeation of the 17-4 PH steel plate [20]. The polished specimen was firstly electroplated by a layer of Pd with the thickness of 0.2 lm. The cathodic side (hydrogen entry) was galvanostatically po-
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larized and charged at a constant current density of 20 mA/cm2 in 0.1 N NaOH solution, which was added with 20 mg As2 O3 /l as a hydrogen adsorption promoter. The anodic side (hydrogen exit) was held at a constant potential of 200 mV vs saturated calomel electrode in 0.1 N NaOH at room temperature. Solutions on both sides of the cells were deoxygenated by continuously purging nitrogen bubbles. The potentiostatic current expressed a direct measurement of the hydrogen flow rate. Experimental permeation transients were recorded on a strip chart recorder. The steady-state hydrogen flux (J1 , mol H/s m2 ) through the specimen is measured in terms of steady-state current density i1 p , and it can be converted to hydrogen permeation flux according to the following equation: J1 ¼ i 1 p =nF
ð2Þ
where n is the number of electrons transferred; and F, the Faraday’s constant. The permeation flux can be defined by: J1 L ¼ i 1 p L=nF
ð3Þ
where L is the specimen thickness. The effective hydrogen diffusivity Deff (m2 /s) can be calculated from the transient permeation curve based on the well-known time lag method: Deff ¼ L2 =6tL
ð4Þ
where tL is the lag time, which is simply defined as 0.63 times the steady-state value. The apparent hydrogen solubility Capp (mol H/m3 ) is determined by: Capp ¼ J1 L=Deff
ð5Þ
3. Results and discussion 3.1. Microstructural observations and microhardness measurements Fig. 2 displays the metallograph of cross-section and microhardness distribution of the LA specimen. Traditional laser-transformation hardening of the carbon steel results in the formation of austenite upon heating cycle, and the austenite transforms into hard martensite during subsequent conduction quench [21]. Therefore, a significant increase in the surface hardness of the steel can be observed. The hardened specimen, in general, consists of hardened zone, partially transformed zone, overtempered zone and heat-unaffected base metal [21]. In contrast to the above process, laser-surface-annealing of aged 17-4 PH stainless steel develops a soft surface layer due to its low carbon content. As shown in Fig. 2, the LA specimen has a lenticular region adjacent to the surface with the microhardness around Hv 348. Furthermore, the microhardness is increased with increasing the distance below the soft surface layer. The distinct microstructures corresponding to various locations below the surface primarily depend on the thermal history at those specific sites during laser treatment. If the heating cycle exceeds the solution temperature of the alloy with
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Fig. 2. A typical (a) metallograph and (b) microhardness distribution of the LA specimen (BM: base metal).
sufficient time period, most precipitates in the alloy will dissolve into the matrix. The inherent low carbon content of the 17-4 PH steel provides little hardening effect upon cooling. Thus, the decrease in microhardness in the surface lenticular region can be accounted for the dissolution of strengthening precipitates. With increasing the distance below specimen surface, the dissipation of thermal energy by conducting heat into the base metal results in decreasing the peak temperature of the thermal cycle. If the peak temperature is not high enough to dissolve the precipitates in the matrix completely, the partially transformed zone will be formed after laser annealing. As demonstrated in Fig. 2(a), the bright-etched region located at the distance between 0.80 and 1.70 mm below the surface is considered as the partially transformed zone. The microhardness in this region mainly depends upon the degree of precipitate dissolution into the matrix. Further increasing the distance toward to the specimen center, the precipitates in the matrix become coarsened after laser treatment, displaying the feature of over-aged microstructure. Finally, the degraded laser energy has no substantial effect on the specimen and exhibits the microstructure of base metal. Fig. 3 shows bright field image of the TEM micrograph for SA 17-4 PH steel. The microstructures of the SA specimen mainly consist of lath martensite with high dislocation density, and there are no precipitates observed in both the martensite matrix and grain boundaries. It was also reported that very fine copper-rich pre-
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Fig. 3. The bright field of the TEM micrograph for the SA specimen.
cipitates are uniformly distributed in the matrix of 17-4 PH steel aged at 499 °C for 2 h [3,22]. Fig. 4 shows the TEM micrograph of the peak-aged (H900) specimen. Its matrix is primarily comprised of tempered lath martensite and fine precipitates. However, no precipitates are observed along the prior austenite grain boundaries as shown in the figure. Fig. 5 displays TEM micrographs of the over-aged (H1025) specimen. Fig. 5(a) and (c) are bright field images of the H1025 specimen. The dark-field image of a (2 0 0) reverted austenite reflection is shown in Fig. 5(b). Its blocky morphology in the matrix is consistent with the published literature [3,22]. Based on the dark-field image of a (2 0 0) e-copper precipitate reflection as demonstrated in Fig. 5(d), the incoherent copper-rich precipitates are extensively located in both the matrix and
Fig. 4. The bright field of the TEM micrograph for the peak-aged (H900) specimen.
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Fig. 5. TEM micrographs of the over-aged (H1025) specimen: (a) the bright field image, (b) the dark-field image of a (2 0 0) reverted austenite reflection, (c) the bright field image, (b) the dark-field image of a (2 0 0) e-copper precipitate reflection.
prior austenite grain boundaries. As displayed in the figure, grain boundaries are decorated with relatively coarse precipitates. Because both the e-Cu precipitate and reverted austenite have the same crystal structure and very close lattice constants below 1% in difference, it is not possible to distinguish them based on TEM structural analysis only. It is reported that a thin layer of austenite along lath boundaries and intralath Cu-rich phase are observed in 17-4 PH steel over-aged at 600 °C for 1 h [3]. Moreover, the Cu-rich precipitates preferentially nucleate and grow along the prior austenite grain boundaries [22]. In addition, the blocky austenite is coarsened at a much higher rate than the Cu-rich precipitates [3,22]. Austenites in the grain interior and along grain boundaries, which are much coarser than incoherent e-Cu precipitates, are observed in the over-aged T250 maraging steel [23]. Based on the above discussion, it can be deduced that the e-Cu phase will be the major precipitates in prior austenite grain boundaries for the over-aged 17-4 PH steels. 3.2. Notch tensile tests Fig. 6 shows the NTS of various specimens tested in air, gaseous hydrogen and saturated H2 S solution, respectively. The peak-aged specimen (H900) exhibits the
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Fig. 6. NTS of various specimens tested in air, gaseous hydrogen and saturated H2 S solution.
highest NTS tested in air among all specimens due to its precipitate-strengthened microstructure. Both the over-aged specimen (H1025) and LA specimen have similar NTS, and their strength is inferior to that of peak-aged specimen. The LA specimen has LAZs on its surfaces, so lower strength than that of the H900 specimen is expected. Similarly, the SA specimen without precipitate-strengthening effect has the lowest NTS among the specimens. As the environment changes from air into gaseous hydrogen, the decrease in NTS for all the specimens is very limited even for the H900 specimen. The result demonstrates that the 17-4 PH stainless steel is resistant to gaseous hydrogen embrittlement. However, it is noted that the decrease of NTS in the saturated H2 S solution is very pronounced for all the specimens. Comparing the NTS tested in air with that in hydrogen-related environments, the NTS loss could be used as an index to assess the hydrogen embrittlement susceptibility among all specimens. In other words, higher NTS loss of the specimen corresponds to higher susceptibility to hydrogen embrittlement. The influence of various environments on the NTS loss is shown in Fig. 7. In gaseous hydrogen, the NTS loss for all specimens is below 6%. As a whole, all specimens have low susceptibility to gaseous hydrogen embrittlement. In addition, the SA specimen has the lowest NTS loss in contrast to H1025 specimen with the highest NTS loss among the specimens in H2 . The lowest NTS loss of the SA specimen in H2 is attributed to its inherent high ductility and low strength. Local yielding of a ductile material can result in notch tip blunting during straining, and the stress-assisted hydrogen diffusion into the notch front is greatly reduced.
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Fig. 7. The influence of hydrogen-related environment on NTS loss for various specimens.
A drastic drop in NTS for the specimens tested in the H2 S solution is observed. It is noted that the SA specimen shows the highest resistance to gaseous hydrogen embrittlement, but has the greatest percentage NTS loss among all specimens tested in the H2 S solution. It demonstrates that the soft SA specimen is highly susceptible to sulfide SCC. On the other hand, the H900 specimen with the highest strength in air shows the lowest susceptibility to sulfide SCC. Furthermore, the LA specimen tested in air has a slightly higher strength than H1025 specimens, but it is changed if tested in H2 S solution. It can be attributed to the presence of LAZs in the LA specimen, in which the material in SA condition has the highest degradation of NTS in H2 S solution. The high NTS loss of LA specimen also confirms that the 17-4 PH stainless steel in the SA condition is very sensitive to sulfide SCC. Therefore, the sulfide SCC resistance of peak-aged 17-4 PH steel cannot be improved by laserannealing treatment. Fig. 8 are SEM fractographs revealing the fracture appearance ahead of the notch front for various specimens tested in gaseous hydrogen. Macroscopic fracture morphology of various specimens consists of two major portions, the flat fractured delta (FD) region at center and two slant fracture (SF) regions in basin shape adjacent to FD region. The flat fracture region represents loading under higher constraint and fractures in plain strain condition during tensile test. Whereas, the near surface regions with SF is deformed in plane stress condition. The higher the strength of the alloy, the wider area of the material will be stressed under plain strain condition. Therefore, higher hydrogen embrittlement susceptibility of the alloy will be expected. The SA specimen with lower NTS in air possesses a lesser extent of FD region than other specimens as shown in Fig. 8. Increasing the strength of the alloy, the
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Fig. 8. SEM fractographs of tensile fractured specimens tested in gaseous hydrogen: (a) SA, (b) H900, (c) LA and (d) H1025 specimens.
relative proportion of slant to flat fracture is changed. In contrast to the SA specimen, the H900 specimen exhibits the largest extent of delta region among the specimens as displayed in Fig. 8(b). It is observed that the delta region of H900 specimen is much flatter than that of the other specimens. Consequently, H900 specimen fails with the trend of brittle fracture, and it will be confirmed in following SEM observation. The presence of soft LAZs on the surfaces of LA specimen will have less plastic constraint of the specimen during tensile deformation. Hence, the tendency of brittle fracture in central portion of the LA specimen is decreased. This feature is proved by the existence of ductile fracture in the delta region of the LA specimen tested in gaseous hydrogen (Fig. 8(c)). The result also demonstrates that the application of laser-surface-annealing treatment on the H900 specimen can reduce the susceptibility of such a specimen to gaseous hydrogen embrittlement. Macroscopic fractographs of notched tensile specimens tested in the H2 S solution are shown in Fig. 9. As mentioned previously, all the specimens suffer from severe hydrogen embrittlement and experience an obvious deterioration of NTS in H2 S solution. The fracture appearance shows SF on the surface for all specimens tested in air or gaseous hydrogen. However, flat fracture instead of SF is found for the specimens suffered from sulfide SCC, as shown in Fig. 9. The fracture surface of the notched specimen is surrounded by embrittled zone (EZ) as displayed in the figure.
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Fig. 9. Macroscopic photographs of tensile fractured specimens tested in H2 S-saturated solution: (a) SA, (b) H900 and (c) LA specimens. The extent of the EZ is delimited by the dashed line.
Wide extent of brittle fracture for the specimens is responsible for high NTS loss in the H2 S solution. It is believed that larger portion of flat fracture region in the fractured surface represents the longer time period available for the hydrogen atoms to induce subcritical crack growth before final failure occurs. Thus, comparing the embrittled area or depth among various specimens can provide a useful index to evaluate the hydrogen embrittlement susceptibility of all specimens tested in H2 S solution. In Fig. 9, both left and right sides of the micrographs are close to the notch front, and the upper and lower sides of the micrographs are free surfaces of the tested specimen. Based on the experimental observation, the depth of EZ ahead of the notch front is deeper than that ahead of free surfaces as shown in Fig. 9(a) and (b). The evidence illustrates that hydrogen atoms tend to diffuse into the highly stressed locations and
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embrittle such regions. It is also noted that the SA specimen experiences a much deeper brittle depth than all other specimens in H2 S solution (Fig. 9(a)), and it is well consistent with the observed high NTS loss of SA specimen tested in H2 S solution. In contrast to SA specimen, H900 specimen has a relatively shallow depth of brittle fracture (Fig. 9(b)), and it is correlated with the relatively low NTS loss in H2 S solution. It is noted that the upper and lower sides of the micrograph shown in Fig. 9(c) are the LAZs. The fractograph of LA specimen also indicates that embrittled depth at the specimen’s surface is much deeper than that at the notch front (Fig. 9(c)). This phenomenon is responsible for the significant increment of NTS loss for LA specimen tested in the H2 S solution. The fracture appearance near crack initiation sites for various specimens tested in specified environments is shown in Fig. 10. The fractographs of various specimens tested in air or gaseous hydrogen shows ductile dimple fracture except for H900 specimen tested in H2 . As shown in Fig. 8(b), the fracture surface of H900 specimen tested in gaseous hydrogen consists of wide flat delta region ahead of notch front. Quasi-cleavage fracture is found in this region demonstrating the evidence of hydrogen embrittlement (Fig. 10(a)). It is worth noting that the presence of brittle fracture feature in H900 specimen tested in gaseous hydrogen does not lead to an
Fig. 10. SEM fractographs illustrating: (a) quasi-cleavage fracture of the H900 specimen in H2 , (b) quasicleavage fracture of the SA specimen in the H2 S solution, (c) intergranular fracture of the H1025 specimen in H2 S and (d) intergranular fracture of the LA specimen tested in H2 S solution at the depth of 1.8 mm below the surface.
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obvious decrease in NTS. Enhancing hydrogen–material interaction by either decreasing the displacement rate or increasing the gaseous hydrogen pressure during testing may result in the occurrence of pronounced hydrogen embrittlement. The severe deterioration of NTS for the specimens tested in H2 S solution is associated with the change of fracture morphology. Both SA and H900 specimens reveal quasicleavage fracture (Fig. 10(b)), and mainly intergranular fracture is found in the H1025 specimen (Fig. 10(c)). Furthermore, intergranular fracture is observed at a distance about 1.8 mm below the surface for tensile fractured LA specimen tested in the H2 S solution (Fig. 10(d)). These will be further discussed in the following paragraphs. 3.3. Hydrogen permeation tests Table 1 lists the results of electrochemical permeation tests including: effective diffusivity (Deff ), permeation flux (J1 L) and apparent hydrogen solubility (Capp ) in SA, H900 and H1025 specimens, respectively. It indicates that the difference in hydrogen permeation flux among various specimens is not significant. However, the effective diffusivity of SA specimen is much higher than that of other specimens. It is about three times greater than that of H900 and H1025 specimens. Meanwhile, the SA specimen has the lowest apparent hydrogen solubility. In other words, hydrogen atoms tend to diffuse at a much faster rate in the SA specimen. On the other hand, H1025 specimen has the lowest effective diffusivity and highest apparent hydrogen solubility among all specimens. Therefore, large amount of hydrogen atoms can be trapped in the H1025 specimen. In general, hydrogen traps in the alloy can be classified as reversible and irreversible trapped sites according to their binding energy [24,25]. The existence of lattice defects such as dislocations, grain boundaries and precipitate/matrix interfaces can trap a large amount of hydrogen atoms. Increasing the number of trapping sites in the alloy, the hydrogen solubility in the alloy is also increased, and it results in decreasing the hydrogen diffusivity. It is reported that the hydrogen diffusivity of martensitic steels lower than that of a-iron is considered to be a multiple trapping effect for hydrogen, from the complex imperfections in the steel such as dislocations, grain boundaries and solute atoms, etc. [26]. In addition, irreversible hydrogen traps can always lead to a greater decrease in hydrogen diffusivity than reversible hydrogen traps [24]. It is reported that irreversible traps result in a lowering hydrogen diffusivity through the material, hence, delaying the dangerous accumulation of hydrogen on Table 1 Results of electrochemical permeation test for 17-4 PH steel with three different heat treatments Properties
SA
H900
H1025
Deff (m2 /s) J1 (mol H/s m) Capp (mol H/m3 )
2:74 1012 58:5 1011 215
0:97 1012 51:5 1011 531
0:81 1012 59:8 1011 738
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critical defect [26]. However, Pound indicates that hydrogen embrittlement susceptibility of the alloy increases with increasing the irreversible trapping constant [27,28]. Parvathavarthini also points out that the lattice defects are the dominant features, which govern the characteristics of hydrogen diffusivity and solubility of 9Cr–1Mo–0.1C steel [29]. As-quenched low carbon lath martensite with the microstructure of high dislocation density and high degree of lattice strain has the highest solubility and lowest diffusivity. For the normalized specimens, the decrease in defect density by using slower cooling rate results in the decrease in trapping sites, and it is correlated with the decrease in solubility and increase in apparent hydrogen diffusivity. The as-quenched 2.25Cr–Mo steel also has the lowest diffusion coefficient; after tempering an increase in the diffusion coefficient is attributed to a reduction in reversible traps [30]. For the 17-4 PH steel, the diffusivity decreases with increasing the aging temperature, which varies in opposite direction of the Cr–Mo steels. It reveals that the variation in hydrogen permeation properties between various materials will be quite unpredictable. The driving force of hydrogen diffusion in the alloy depends not only on the concentration gradient but also on stress gradient. The degree of embrittlement for notch specimen predominantly is governed by surface hydrogen absorption and transport of hydrogen atoms into the highly stressed region. The transportation of hydrogen into highly stressed sites can be performed by lattice diffusion and/or dislocation sweeping mechanism [31,32]. For alloys with the microstructure, in which hydrogen atoms can diffuse fast and be trapped less, the accumulation of embrittling species in the critical stressed location of the alloy is much more effective. Therefore, a severe degradation of NTS and deep EZ will be expected. The microstructure of SA specimen is mainly comprised of lath martensite with high dislocation density. Large amount of reversible trap sites and relatively few strong irreversible trapping sites are available in the specimen. This feature is consistent with the permeation test result, i.e. highest hydrogen diffusivity and lowest apparent hydrogen solubility found in the SA specimen. As discussed previously, the SA specimen shows the highest susceptibility to sulfide SCC, i.e. the highest NTS loss in H2 S solution. The sufficient supplement of embrittling species deteriorates the specimen effectively, and the wide EZ is also consistent with such characteristics. Meanwhile, the uniformly trapped hydrogen predominantly in the grain interior promotes quasi-cleavage fracture of the SA specimen tested in the H2 S solution. There are many trap sites in both H900 and H1025 specimens. The inward diffusion of hydrogen is highly interfered by their microstructural inhomogenities, e.g. the distribution of both Cu-rich precipitates and interlath austenite. Thus, both specimens have much lower Deff and higher Capp than those of the SA specimen do. Retained austenite and/or the interface between retained austenite and matrix act as strong trapping sites for hydrogen, and they can also dissolve large amount of hydrogen [26,33]. Thus, the difference in permeation properties between H900 and H1025 specimens can be partly attributed to the presence of reverted austenite in the H1025 specimen. Grain boundaries of H1025 specimens are decorated with many precipitates, as shown in Fig. 5(d). Hydrogen can diffuse and be trapped along the grain boundaries. In contrast, the grain interior with blocky reverted austenite and
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incoherent precipitates has higher solubility to dissolve large amount of hydrogen. Therefore, the existence of hydrogen reservoir in grains can effectively delay the time period required to induce hydrogen embrittlement of gain interior. Hydrogen trapped along grain boundary results in decreasing the cohesive strength between grains, so the fracture of H1025 specimen in the H2 S solution is primarily along prior austenite grain boundaries. As shown in Fig. 10(d), intergranular fracture of the LA specimen is found in the depth of 1.8 mm below the surface. According to the microhardness–depth distribution shown in Fig. 2(b), the above region can be classified as the over-aged zone of the LA specimen. The over-aged precipitates in LA specimen may offer hydrogentrapping sites similar to those of H1025 specimens. Therefore, intergranular fracture of the LA specimen in depth of 1.8 mm below the surface is observed. The H900 specimen has a less extent of EZ and associates with a relatively low NTS loss in the H2 S solution as compares with other specimens. The microstructure of peak-aged 17-4 PH steel consists of a uniformly distributed fine precipitates in the matrix. It is deduced that hydrogen atoms are trapped mainly in the particles/matrix interfaces. The uniformly trapped hydrogen atoms in the matrix of H900 specimen account for its superior NTS in comparison with all other specimens. Pressouyre reports that the microstructure of the Fe–Ti alloy consists of a homogeneous and fine distribution of reversible hydrogen traps [34,35]. The presence of uniform reversible traps, which distribute the hydrogen innocuously, will reduce the extent of hydrogen-induced cracking [34,35]. The 17-4 PH stainless steel in peak-aged condition behaves the similar characteristics. It is inferred that decohesion of precipitate/ matrix interfaces in the matrix results in the formation of quasi-cleavage fracture of the H900 specimen tested in the H2 S solution.
4. Conclusions 1. Regardless of testing environments, the peak-aged (H900) specimen demonstrates the highest NTS among all specimens. All specimens display low NTS loss in gaseous hydrogen but significantly suffer from sulfide SCC, especially for the SA specimen. 2. Hydrogen permeation properties of various specimens are strongly related to the microstructure of 17-4 PH stainless steel. Both high hydrogen diffusivity and low hydrogen solubility of the SA specimen provide the mechanism of rapid transportation of hydrogen atoms to the highly stressed region with little interference. Therefore, severely deteriorated NTS is observed for the SA specimen tested in the H2 S solution. 3. Microstructural observation of the over-aged (H1025) specimen reveals those blocky austenites together with Cu-rich precipitates uniformly distributed in grain interior. Dense and coarse precipitates are delineated along prior austenite grain boundaries. The H1025 specimen with the lowest Deff and highest Capp among all specimens can be attributed to the presence of numerous trapping sites prohibiting the inward diffusion of hydrogen atoms. Grain interior has higher resistance to
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hydrogen embrittlement than grain boundaries. Hydrogen atoms are trapped along grain boundaries and lead to the formation of intergranular fracture for H1025 specimen tested in the H2 S solution. 4. Fine and homogeneously distributed precipitates in the H900 specimen account for its superior NTS tested in the H2 S solution. Hydrogen atoms are trapped predominantly in the precipitate/matrix interfaces. The decohesion of these interfaces results in the formation of quasi-cleavage fracture of the H900 specimen tested in H2 S solution. 5. Because the microstructure in the SA condition is very susceptible to hydrogen embrittlement in the H2 S solution, the application of laser-surface-annealing treatment on the peak-aged specimen (LA specimen) cannot improve its resistance to sulfide SCC.
Acknowledgements The authors gratefully acknowledge the financial support from the National Science Council of ROC (contract no. NSC 89-2216-E-019-001).
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B.G. Pound, Corrosion 45 (1989) 18. N. Parvathavarthini, S. Saroja, R.K. Dayal, J. Nucl. Mater. 264 (1999) 35. R. Valentini, A. Solina, Mater. Sci. Tech. Ser. 10 (1994) 908. J.K. Tien, A.W. Thompson, I.M. Bernstein, R.J. Richards, Metall. Mater. Trans. 7A (1976) 821. T. Tabata, H.K. Birnbaum, Scripta Mater. 17 (1983) 947. S.L.I. Chan, H.L. Lee, J.R. Yang, Metall. Mater. Trans. 22A (1991) 2579. G.M. Pressouyre, I.M. Bernstein, Acta Metall. Mater. 27 (1979) 89. G.M. Pressouyre, Acta Metall. Mater. 28 (1980) 895.
Notch tensile properties of laser-surface-annealed 17-4 PH stainless steel in hydrogen-related environments L.W. Tsay a
a,*
, W.C. Lee a, R.K. Shiue b, J.K. Wu
a
Institute of Materials Engineering, National Taiwan Ocean University, Pei-Ning Road, Keelung 202, Taiwan, ROC b Department of Materials Science and Engineering, National Dong Hwa University, Hualien 974, Taiwan, ROC Received 21 September 2001; accepted 27 December 2001
Abstract Slow displacement rate tensile tests were performed to determine the notched tensile strength (NTS) of 17-4 PH stainless steel with various microstructures in hydrogen-related environments. Solution-annealed (SA), peak-aged (H900), over-aged (H1025), and laserannealed (LA) specimens were included in the study. Based on the results of NTS in air, the NTS loss in both gaseous hydrogen and H2 S-saturated solution was used to access the detrimental effects of hydrogen in 17-4 PH steel subjected to different treatments. Electrochemical permeation tests were also employed to determine the hydrogen permeation characteristics of the 17-4 PH steel plate with various microstructures. The result indicates that all the specimens have low NTS loss in gaseous hydrogen but significantly suffer from sulfide stress corrosion cracking (SCC), especially for the soft SA specimen. It was deduced that high hydrogen diffusivity and less trapped hydrogen atoms in the SA matrix provided rapid transport of massive hydrogen atoms into highly stressed region, and deteriorated the NTS tested in the saturated sulfide solution. On the other hand, H1025 specimen consists of the blocky austenite together with Cu-rich precipitates uniformly distributed in the grain interior; dense and coarse precipitates are also observed along prior austenite grain boundaries. Hydrogen atoms tend to be trapped along grain boundaries, and lead to the formation of intergranular fracture for H1025 specimen tested in the H2 S solution. Fine and homogeneously distributed precipitates in the H900 matrix result in uniformly trapping of hydrogen atoms, so it behaves superior properties than other specimens. The decohesion of precipitate/matrix interfaces induces quasi-cleavage fracture of the H900 specimen tested in H2 S solution. Finally, the application
*
Corresponding author. Tel.: +886-2-24622192; fax: +886-2-24625324. E-mail addresses: [email protected] (L.W. Tsay), [email protected] (R.K. Shiue).
0010-938X/02/$ - see front matter Ó 2002 Elsevier Science Ltd. All rights reserved. PII: S 0 0 1 0 - 9 3 8 X ( 0 2 ) 0 0 0 2 3 - 9
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of laser-annealing treatment on the H900 specimen cannot improve its resistance to sulfide SCC, because the laser-annealed zone is susceptible to hydrogen embrittlement in the H2 S solution. Ó 2002 Elsevier Science Ltd. All rights reserved. Keywords: C. Hydrogen embrittlement; Sulfide stress corrosion cracking; 17-4 PH stainless steels; Notch tensile strength
1. Introduction The 17-4 PH (AISI 630) steel is a precipitation-hardened martensitic stainless steel with high strength as well as moderate corrosion resistance, and it has been widely used in various industries [1,2]. After solution-annealed (SA) treatment, 17-4 PH stainless steel can be hardened by various aging treatments. The formation of coherent copper-rich clusters in the peak-aged condition is primarily responsible for the rapid increase of its hardness [3]. As the aging temperature is further increased above 600 °C, both the formation of incoherent e precipitates in the matrix and a significant amount of martensite transform into austenite along martensite lath boundaries [3]. The precipitation of incoherent e phase during over-aging will have depleted substantial amounts of Cu and Ni from the matrix, resulting in enhanced electrochemical dissolution of the martensite matrix [4]. It is well known that surface properties of many structural alloys can be greatly improved by the laser surface treatment [5,6]. After irradiating by the laser beam, the age-hardened alloy forms a soft ductile surface region in SA condition, which is socalled laser-annealed zone (LAZ) [7]. The composite region is characterized by such soft LAZs on its outer surfaces and the untransformed base metal sandwiched between two LAZs [7]. Meanwhile, it is reported that the improvement of fatigue crack growth is very pronounced in the region preceding the composite region for peakaged 17-4 PH stainless steel subjected to laser-surface-annealed treatment [8]. The stress corrosion cracking (SCC) of an alloy is strongly related to three important factors, including the microstructure of the alloy, tri-axial stress state of the specimen and environment of the alloy therein. In general, the hardened low alloy steel applied in hydrogen-related environments under high tensile stress is much more susceptible to SCC than the highly tempered one employed in the same environment under compressive stress. Traditionally, the hydrogen embrittlement susceptibility of high strength alloys can be greatly improved by decreasing the yield strength of the alloy. The age-hardened alloy in over-aged condition is usually more resistant to hydrogen embrittlement than that in peak-aged condition. 17-4 PH stainless steel is extensively used in the power industry. In geothermal power applications, the presence of H2 S in the environment may result in severe hydrogen embrittlement of the alloy. It is reported that both hydrogen embrittlement and anodic dissolution mechanism are responsible for SCC of the 17-4 PH stainless steel in an acidic NaCl solution [9]. The 17-4 PH stainless steel in the SA state is found to be very resistant to SCC in anodic potential environment [10]. Thus, it was deduced that the peak-aged 17-4 PH stainless steel with LAZs on the surfaces might
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not only possess adequate strength but also exhibited improved resistance to hydrogen embrittlement. Fractographs of many low alloy steels susceptible to hydrogen embrittlement depends on its microstructure, crack-tip stress intensity and hydrogen concentration around the crack tip [11]. Moreover, the susceptible alloy suffered from hydrogen embrittlement usually associates with the change of fracture modes [12,13]. SCC of low strength steels is characterized by transgranular fracture in contrast to intergranular fracture of high strength ones [12–14]. It has also been pointed out that hydrogen atoms are prone to diffuse into locations with high hydrostatic stress [15– 17]. In addition, hydrogen atoms tend to be accumulated and embrittle the elastic– plastic boundary of the high strength steel [17]. Consequently, notched tensile test under a slow displacement rate is a suitable experimental process in evaluating the detrimental effect of hydrogen to the alloy with various microstructures. The purpose of this investigation was focused on the influence of hydrogen embrittlement on 17-4 PH steel with various microstructures. Both degradation of notched tensile strength (NTS) and related fracture characteristics of 17-4 PH stainless steel were studied extensively. Slow displacement rate tensile tests were performed to enhance the interaction between the alloy and the hydrogen-related environment. The detrimental effect of hydrogen was accessed by comparing the NTS in air, gaseous hydrogen and H2 S-saturated NACE solution, respectively. Electrochemical permeation was employed to determine hydrogen permeation properties of the specimens aged at different conditions. Fractographic observations were carried out on the tensile fractured specimen focused on the neighborhood of the crack initiation sites, in which the changes in fracture morphology were usually observed. The deterioration of NTS was strongly related to the extent of embrittled area on its fracture surface. The influence of hydrogen embrittlement on the NTS of the specimens was explained on the basis of hydrogen/microstructure interactions, i.e. in terms of hydrogen diffusivity and hydrogen solubility of the alloy.
2. Experimental procedures The chemical composition in wt.% of the 17-4 PH stainless steel was 15.71Cr, 4.35Ni, 3.46Cu, 0.70Si, 0.04C, 0.38Mn, 0.025P, 0.013S and balance Fe. The asreceived steel plate with 5.6 mm in thickness was ground into 5.0 mm thick. It was firstly austenized at 1038 °C for 1 h and subsequently cooled by purging highpressure nitrogen; it was specified as the SA specimen with the hardness of HRC 32. The following aging treatment was performed at 482 °C for 1 h (H900) and 552 °C for 4 h (H1025), respectively. The hardness of peak-aged specimen (H900) was HRC 42, and the hardness of over-aged specimen (H1025) was HRC 34. Laser surface treatments were carried out by using a Rofin-Sinar 5 kW CO2 laser. The size of rectangular laser beam on the specimen was about 6 mm 25 mm with uniform energy intensity within the irradiated area. In order to enhance the absorptivity of laser irradiated surface, the specimen was sprayed by carbon black paint prior to laser treatment. Some H900 specimens were subjected to LA treatment.
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Fig. 1. Schematic diagram showing the dimensions of notched tensile specimen used in the experiment.
After laser irradiation, the specimen comprised of the composite region was recognized as LA specimens. The laser energy used in the experiment was 2000 W, and the scan rate of laser beam was set at 600 mm/min throughout the experiment. The microhardness measurement was performed along the centerline of the LA specimen. The schematic dimensions of notched tensile specimen used in this research are shown in Fig. 1. It is noted that the central shaded area in the figure is the LAZ. The notch tip radius shown in Fig. 1 is about 100 lm. Initial tensile tests were performed in air with the controlled crosshead speed of 0.75 mm/min. The effect of hydrogen embrittlement on the deterioration of NTS was evaluated by installing the specimen into a Teflon chamber with controlled atmosphere, i.e., gaseous hydrogen at 2 atm or saturated H2 S solution. It was reported that sulfide SCC was a manifestation of hydrogen embrittlement [18,19]. The H2 S solution was prepared according to NACE standard (TM-01-77-86). The solution was purged with nitrogen gas for at least 1 h to remove O2 before testing, and H2 S was continuously bubbled into the solution to ensure the solution properly saturated. In order to clarify the effect of hydrogen embrittlement on the alloy, a much slower crosshead speed of 7:5 103 mm/min was chosen during the tensile test. All NTS data were averaged from at least three specimens for each test condition. The degree of susceptibility to hydrogen embrittlement of various specimens can be expressed in the percentage loss of NTS under hydrogen charging as given below: NTS loss ð%Þ ¼
NTS ðin airÞ NTS ðin H2 or H2 SÞ 100% NTS ðin airÞ
ð1Þ
Tensile fractured appearance of various specimens was examined by a Hitachi S4100 field emission scanning electron microscope (SEM). Thin foil specimens for transmission electron microscope (TEM) observation were examined by a JEOL2000EX microscope with the operation voltage of 200 kV. Electrochemical permeation technique originally developed by Devanathan and Stachurski was employed to determine the hydrogen permeation of the 17-4 PH steel plate [20]. The polished specimen was firstly electroplated by a layer of Pd with the thickness of 0.2 lm. The cathodic side (hydrogen entry) was galvanostatically po-
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larized and charged at a constant current density of 20 mA/cm2 in 0.1 N NaOH solution, which was added with 20 mg As2 O3 /l as a hydrogen adsorption promoter. The anodic side (hydrogen exit) was held at a constant potential of 200 mV vs saturated calomel electrode in 0.1 N NaOH at room temperature. Solutions on both sides of the cells were deoxygenated by continuously purging nitrogen bubbles. The potentiostatic current expressed a direct measurement of the hydrogen flow rate. Experimental permeation transients were recorded on a strip chart recorder. The steady-state hydrogen flux (J1 , mol H/s m2 ) through the specimen is measured in terms of steady-state current density i1 p , and it can be converted to hydrogen permeation flux according to the following equation: J1 ¼ i 1 p =nF
ð2Þ
where n is the number of electrons transferred; and F, the Faraday’s constant. The permeation flux can be defined by: J1 L ¼ i 1 p L=nF
ð3Þ
where L is the specimen thickness. The effective hydrogen diffusivity Deff (m2 /s) can be calculated from the transient permeation curve based on the well-known time lag method: Deff ¼ L2 =6tL
ð4Þ
where tL is the lag time, which is simply defined as 0.63 times the steady-state value. The apparent hydrogen solubility Capp (mol H/m3 ) is determined by: Capp ¼ J1 L=Deff
ð5Þ
3. Results and discussion 3.1. Microstructural observations and microhardness measurements Fig. 2 displays the metallograph of cross-section and microhardness distribution of the LA specimen. Traditional laser-transformation hardening of the carbon steel results in the formation of austenite upon heating cycle, and the austenite transforms into hard martensite during subsequent conduction quench [21]. Therefore, a significant increase in the surface hardness of the steel can be observed. The hardened specimen, in general, consists of hardened zone, partially transformed zone, overtempered zone and heat-unaffected base metal [21]. In contrast to the above process, laser-surface-annealing of aged 17-4 PH stainless steel develops a soft surface layer due to its low carbon content. As shown in Fig. 2, the LA specimen has a lenticular region adjacent to the surface with the microhardness around Hv 348. Furthermore, the microhardness is increased with increasing the distance below the soft surface layer. The distinct microstructures corresponding to various locations below the surface primarily depend on the thermal history at those specific sites during laser treatment. If the heating cycle exceeds the solution temperature of the alloy with
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Fig. 2. A typical (a) metallograph and (b) microhardness distribution of the LA specimen (BM: base metal).
sufficient time period, most precipitates in the alloy will dissolve into the matrix. The inherent low carbon content of the 17-4 PH steel provides little hardening effect upon cooling. Thus, the decrease in microhardness in the surface lenticular region can be accounted for the dissolution of strengthening precipitates. With increasing the distance below specimen surface, the dissipation of thermal energy by conducting heat into the base metal results in decreasing the peak temperature of the thermal cycle. If the peak temperature is not high enough to dissolve the precipitates in the matrix completely, the partially transformed zone will be formed after laser annealing. As demonstrated in Fig. 2(a), the bright-etched region located at the distance between 0.80 and 1.70 mm below the surface is considered as the partially transformed zone. The microhardness in this region mainly depends upon the degree of precipitate dissolution into the matrix. Further increasing the distance toward to the specimen center, the precipitates in the matrix become coarsened after laser treatment, displaying the feature of over-aged microstructure. Finally, the degraded laser energy has no substantial effect on the specimen and exhibits the microstructure of base metal. Fig. 3 shows bright field image of the TEM micrograph for SA 17-4 PH steel. The microstructures of the SA specimen mainly consist of lath martensite with high dislocation density, and there are no precipitates observed in both the martensite matrix and grain boundaries. It was also reported that very fine copper-rich pre-
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Fig. 3. The bright field of the TEM micrograph for the SA specimen.
cipitates are uniformly distributed in the matrix of 17-4 PH steel aged at 499 °C for 2 h [3,22]. Fig. 4 shows the TEM micrograph of the peak-aged (H900) specimen. Its matrix is primarily comprised of tempered lath martensite and fine precipitates. However, no precipitates are observed along the prior austenite grain boundaries as shown in the figure. Fig. 5 displays TEM micrographs of the over-aged (H1025) specimen. Fig. 5(a) and (c) are bright field images of the H1025 specimen. The dark-field image of a (2 0 0) reverted austenite reflection is shown in Fig. 5(b). Its blocky morphology in the matrix is consistent with the published literature [3,22]. Based on the dark-field image of a (2 0 0) e-copper precipitate reflection as demonstrated in Fig. 5(d), the incoherent copper-rich precipitates are extensively located in both the matrix and
Fig. 4. The bright field of the TEM micrograph for the peak-aged (H900) specimen.
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Fig. 5. TEM micrographs of the over-aged (H1025) specimen: (a) the bright field image, (b) the dark-field image of a (2 0 0) reverted austenite reflection, (c) the bright field image, (b) the dark-field image of a (2 0 0) e-copper precipitate reflection.
prior austenite grain boundaries. As displayed in the figure, grain boundaries are decorated with relatively coarse precipitates. Because both the e-Cu precipitate and reverted austenite have the same crystal structure and very close lattice constants below 1% in difference, it is not possible to distinguish them based on TEM structural analysis only. It is reported that a thin layer of austenite along lath boundaries and intralath Cu-rich phase are observed in 17-4 PH steel over-aged at 600 °C for 1 h [3]. Moreover, the Cu-rich precipitates preferentially nucleate and grow along the prior austenite grain boundaries [22]. In addition, the blocky austenite is coarsened at a much higher rate than the Cu-rich precipitates [3,22]. Austenites in the grain interior and along grain boundaries, which are much coarser than incoherent e-Cu precipitates, are observed in the over-aged T250 maraging steel [23]. Based on the above discussion, it can be deduced that the e-Cu phase will be the major precipitates in prior austenite grain boundaries for the over-aged 17-4 PH steels. 3.2. Notch tensile tests Fig. 6 shows the NTS of various specimens tested in air, gaseous hydrogen and saturated H2 S solution, respectively. The peak-aged specimen (H900) exhibits the
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Fig. 6. NTS of various specimens tested in air, gaseous hydrogen and saturated H2 S solution.
highest NTS tested in air among all specimens due to its precipitate-strengthened microstructure. Both the over-aged specimen (H1025) and LA specimen have similar NTS, and their strength is inferior to that of peak-aged specimen. The LA specimen has LAZs on its surfaces, so lower strength than that of the H900 specimen is expected. Similarly, the SA specimen without precipitate-strengthening effect has the lowest NTS among the specimens. As the environment changes from air into gaseous hydrogen, the decrease in NTS for all the specimens is very limited even for the H900 specimen. The result demonstrates that the 17-4 PH stainless steel is resistant to gaseous hydrogen embrittlement. However, it is noted that the decrease of NTS in the saturated H2 S solution is very pronounced for all the specimens. Comparing the NTS tested in air with that in hydrogen-related environments, the NTS loss could be used as an index to assess the hydrogen embrittlement susceptibility among all specimens. In other words, higher NTS loss of the specimen corresponds to higher susceptibility to hydrogen embrittlement. The influence of various environments on the NTS loss is shown in Fig. 7. In gaseous hydrogen, the NTS loss for all specimens is below 6%. As a whole, all specimens have low susceptibility to gaseous hydrogen embrittlement. In addition, the SA specimen has the lowest NTS loss in contrast to H1025 specimen with the highest NTS loss among the specimens in H2 . The lowest NTS loss of the SA specimen in H2 is attributed to its inherent high ductility and low strength. Local yielding of a ductile material can result in notch tip blunting during straining, and the stress-assisted hydrogen diffusion into the notch front is greatly reduced.
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Fig. 7. The influence of hydrogen-related environment on NTS loss for various specimens.
A drastic drop in NTS for the specimens tested in the H2 S solution is observed. It is noted that the SA specimen shows the highest resistance to gaseous hydrogen embrittlement, but has the greatest percentage NTS loss among all specimens tested in the H2 S solution. It demonstrates that the soft SA specimen is highly susceptible to sulfide SCC. On the other hand, the H900 specimen with the highest strength in air shows the lowest susceptibility to sulfide SCC. Furthermore, the LA specimen tested in air has a slightly higher strength than H1025 specimens, but it is changed if tested in H2 S solution. It can be attributed to the presence of LAZs in the LA specimen, in which the material in SA condition has the highest degradation of NTS in H2 S solution. The high NTS loss of LA specimen also confirms that the 17-4 PH stainless steel in the SA condition is very sensitive to sulfide SCC. Therefore, the sulfide SCC resistance of peak-aged 17-4 PH steel cannot be improved by laserannealing treatment. Fig. 8 are SEM fractographs revealing the fracture appearance ahead of the notch front for various specimens tested in gaseous hydrogen. Macroscopic fracture morphology of various specimens consists of two major portions, the flat fractured delta (FD) region at center and two slant fracture (SF) regions in basin shape adjacent to FD region. The flat fracture region represents loading under higher constraint and fractures in plain strain condition during tensile test. Whereas, the near surface regions with SF is deformed in plane stress condition. The higher the strength of the alloy, the wider area of the material will be stressed under plain strain condition. Therefore, higher hydrogen embrittlement susceptibility of the alloy will be expected. The SA specimen with lower NTS in air possesses a lesser extent of FD region than other specimens as shown in Fig. 8. Increasing the strength of the alloy, the
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Fig. 8. SEM fractographs of tensile fractured specimens tested in gaseous hydrogen: (a) SA, (b) H900, (c) LA and (d) H1025 specimens.
relative proportion of slant to flat fracture is changed. In contrast to the SA specimen, the H900 specimen exhibits the largest extent of delta region among the specimens as displayed in Fig. 8(b). It is observed that the delta region of H900 specimen is much flatter than that of the other specimens. Consequently, H900 specimen fails with the trend of brittle fracture, and it will be confirmed in following SEM observation. The presence of soft LAZs on the surfaces of LA specimen will have less plastic constraint of the specimen during tensile deformation. Hence, the tendency of brittle fracture in central portion of the LA specimen is decreased. This feature is proved by the existence of ductile fracture in the delta region of the LA specimen tested in gaseous hydrogen (Fig. 8(c)). The result also demonstrates that the application of laser-surface-annealing treatment on the H900 specimen can reduce the susceptibility of such a specimen to gaseous hydrogen embrittlement. Macroscopic fractographs of notched tensile specimens tested in the H2 S solution are shown in Fig. 9. As mentioned previously, all the specimens suffer from severe hydrogen embrittlement and experience an obvious deterioration of NTS in H2 S solution. The fracture appearance shows SF on the surface for all specimens tested in air or gaseous hydrogen. However, flat fracture instead of SF is found for the specimens suffered from sulfide SCC, as shown in Fig. 9. The fracture surface of the notched specimen is surrounded by embrittled zone (EZ) as displayed in the figure.
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Fig. 9. Macroscopic photographs of tensile fractured specimens tested in H2 S-saturated solution: (a) SA, (b) H900 and (c) LA specimens. The extent of the EZ is delimited by the dashed line.
Wide extent of brittle fracture for the specimens is responsible for high NTS loss in the H2 S solution. It is believed that larger portion of flat fracture region in the fractured surface represents the longer time period available for the hydrogen atoms to induce subcritical crack growth before final failure occurs. Thus, comparing the embrittled area or depth among various specimens can provide a useful index to evaluate the hydrogen embrittlement susceptibility of all specimens tested in H2 S solution. In Fig. 9, both left and right sides of the micrographs are close to the notch front, and the upper and lower sides of the micrographs are free surfaces of the tested specimen. Based on the experimental observation, the depth of EZ ahead of the notch front is deeper than that ahead of free surfaces as shown in Fig. 9(a) and (b). The evidence illustrates that hydrogen atoms tend to diffuse into the highly stressed locations and
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embrittle such regions. It is also noted that the SA specimen experiences a much deeper brittle depth than all other specimens in H2 S solution (Fig. 9(a)), and it is well consistent with the observed high NTS loss of SA specimen tested in H2 S solution. In contrast to SA specimen, H900 specimen has a relatively shallow depth of brittle fracture (Fig. 9(b)), and it is correlated with the relatively low NTS loss in H2 S solution. It is noted that the upper and lower sides of the micrograph shown in Fig. 9(c) are the LAZs. The fractograph of LA specimen also indicates that embrittled depth at the specimen’s surface is much deeper than that at the notch front (Fig. 9(c)). This phenomenon is responsible for the significant increment of NTS loss for LA specimen tested in the H2 S solution. The fracture appearance near crack initiation sites for various specimens tested in specified environments is shown in Fig. 10. The fractographs of various specimens tested in air or gaseous hydrogen shows ductile dimple fracture except for H900 specimen tested in H2 . As shown in Fig. 8(b), the fracture surface of H900 specimen tested in gaseous hydrogen consists of wide flat delta region ahead of notch front. Quasi-cleavage fracture is found in this region demonstrating the evidence of hydrogen embrittlement (Fig. 10(a)). It is worth noting that the presence of brittle fracture feature in H900 specimen tested in gaseous hydrogen does not lead to an
Fig. 10. SEM fractographs illustrating: (a) quasi-cleavage fracture of the H900 specimen in H2 , (b) quasicleavage fracture of the SA specimen in the H2 S solution, (c) intergranular fracture of the H1025 specimen in H2 S and (d) intergranular fracture of the LA specimen tested in H2 S solution at the depth of 1.8 mm below the surface.
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obvious decrease in NTS. Enhancing hydrogen–material interaction by either decreasing the displacement rate or increasing the gaseous hydrogen pressure during testing may result in the occurrence of pronounced hydrogen embrittlement. The severe deterioration of NTS for the specimens tested in H2 S solution is associated with the change of fracture morphology. Both SA and H900 specimens reveal quasicleavage fracture (Fig. 10(b)), and mainly intergranular fracture is found in the H1025 specimen (Fig. 10(c)). Furthermore, intergranular fracture is observed at a distance about 1.8 mm below the surface for tensile fractured LA specimen tested in the H2 S solution (Fig. 10(d)). These will be further discussed in the following paragraphs. 3.3. Hydrogen permeation tests Table 1 lists the results of electrochemical permeation tests including: effective diffusivity (Deff ), permeation flux (J1 L) and apparent hydrogen solubility (Capp ) in SA, H900 and H1025 specimens, respectively. It indicates that the difference in hydrogen permeation flux among various specimens is not significant. However, the effective diffusivity of SA specimen is much higher than that of other specimens. It is about three times greater than that of H900 and H1025 specimens. Meanwhile, the SA specimen has the lowest apparent hydrogen solubility. In other words, hydrogen atoms tend to diffuse at a much faster rate in the SA specimen. On the other hand, H1025 specimen has the lowest effective diffusivity and highest apparent hydrogen solubility among all specimens. Therefore, large amount of hydrogen atoms can be trapped in the H1025 specimen. In general, hydrogen traps in the alloy can be classified as reversible and irreversible trapped sites according to their binding energy [24,25]. The existence of lattice defects such as dislocations, grain boundaries and precipitate/matrix interfaces can trap a large amount of hydrogen atoms. Increasing the number of trapping sites in the alloy, the hydrogen solubility in the alloy is also increased, and it results in decreasing the hydrogen diffusivity. It is reported that the hydrogen diffusivity of martensitic steels lower than that of a-iron is considered to be a multiple trapping effect for hydrogen, from the complex imperfections in the steel such as dislocations, grain boundaries and solute atoms, etc. [26]. In addition, irreversible hydrogen traps can always lead to a greater decrease in hydrogen diffusivity than reversible hydrogen traps [24]. It is reported that irreversible traps result in a lowering hydrogen diffusivity through the material, hence, delaying the dangerous accumulation of hydrogen on Table 1 Results of electrochemical permeation test for 17-4 PH steel with three different heat treatments Properties
SA
H900
H1025
Deff (m2 /s) J1 (mol H/s m) Capp (mol H/m3 )
2:74 1012 58:5 1011 215
0:97 1012 51:5 1011 531
0:81 1012 59:8 1011 738
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critical defect [26]. However, Pound indicates that hydrogen embrittlement susceptibility of the alloy increases with increasing the irreversible trapping constant [27,28]. Parvathavarthini also points out that the lattice defects are the dominant features, which govern the characteristics of hydrogen diffusivity and solubility of 9Cr–1Mo–0.1C steel [29]. As-quenched low carbon lath martensite with the microstructure of high dislocation density and high degree of lattice strain has the highest solubility and lowest diffusivity. For the normalized specimens, the decrease in defect density by using slower cooling rate results in the decrease in trapping sites, and it is correlated with the decrease in solubility and increase in apparent hydrogen diffusivity. The as-quenched 2.25Cr–Mo steel also has the lowest diffusion coefficient; after tempering an increase in the diffusion coefficient is attributed to a reduction in reversible traps [30]. For the 17-4 PH steel, the diffusivity decreases with increasing the aging temperature, which varies in opposite direction of the Cr–Mo steels. It reveals that the variation in hydrogen permeation properties between various materials will be quite unpredictable. The driving force of hydrogen diffusion in the alloy depends not only on the concentration gradient but also on stress gradient. The degree of embrittlement for notch specimen predominantly is governed by surface hydrogen absorption and transport of hydrogen atoms into the highly stressed region. The transportation of hydrogen into highly stressed sites can be performed by lattice diffusion and/or dislocation sweeping mechanism [31,32]. For alloys with the microstructure, in which hydrogen atoms can diffuse fast and be trapped less, the accumulation of embrittling species in the critical stressed location of the alloy is much more effective. Therefore, a severe degradation of NTS and deep EZ will be expected. The microstructure of SA specimen is mainly comprised of lath martensite with high dislocation density. Large amount of reversible trap sites and relatively few strong irreversible trapping sites are available in the specimen. This feature is consistent with the permeation test result, i.e. highest hydrogen diffusivity and lowest apparent hydrogen solubility found in the SA specimen. As discussed previously, the SA specimen shows the highest susceptibility to sulfide SCC, i.e. the highest NTS loss in H2 S solution. The sufficient supplement of embrittling species deteriorates the specimen effectively, and the wide EZ is also consistent with such characteristics. Meanwhile, the uniformly trapped hydrogen predominantly in the grain interior promotes quasi-cleavage fracture of the SA specimen tested in the H2 S solution. There are many trap sites in both H900 and H1025 specimens. The inward diffusion of hydrogen is highly interfered by their microstructural inhomogenities, e.g. the distribution of both Cu-rich precipitates and interlath austenite. Thus, both specimens have much lower Deff and higher Capp than those of the SA specimen do. Retained austenite and/or the interface between retained austenite and matrix act as strong trapping sites for hydrogen, and they can also dissolve large amount of hydrogen [26,33]. Thus, the difference in permeation properties between H900 and H1025 specimens can be partly attributed to the presence of reverted austenite in the H1025 specimen. Grain boundaries of H1025 specimens are decorated with many precipitates, as shown in Fig. 5(d). Hydrogen can diffuse and be trapped along the grain boundaries. In contrast, the grain interior with blocky reverted austenite and
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incoherent precipitates has higher solubility to dissolve large amount of hydrogen. Therefore, the existence of hydrogen reservoir in grains can effectively delay the time period required to induce hydrogen embrittlement of gain interior. Hydrogen trapped along grain boundary results in decreasing the cohesive strength between grains, so the fracture of H1025 specimen in the H2 S solution is primarily along prior austenite grain boundaries. As shown in Fig. 10(d), intergranular fracture of the LA specimen is found in the depth of 1.8 mm below the surface. According to the microhardness–depth distribution shown in Fig. 2(b), the above region can be classified as the over-aged zone of the LA specimen. The over-aged precipitates in LA specimen may offer hydrogentrapping sites similar to those of H1025 specimens. Therefore, intergranular fracture of the LA specimen in depth of 1.8 mm below the surface is observed. The H900 specimen has a less extent of EZ and associates with a relatively low NTS loss in the H2 S solution as compares with other specimens. The microstructure of peak-aged 17-4 PH steel consists of a uniformly distributed fine precipitates in the matrix. It is deduced that hydrogen atoms are trapped mainly in the particles/matrix interfaces. The uniformly trapped hydrogen atoms in the matrix of H900 specimen account for its superior NTS in comparison with all other specimens. Pressouyre reports that the microstructure of the Fe–Ti alloy consists of a homogeneous and fine distribution of reversible hydrogen traps [34,35]. The presence of uniform reversible traps, which distribute the hydrogen innocuously, will reduce the extent of hydrogen-induced cracking [34,35]. The 17-4 PH stainless steel in peak-aged condition behaves the similar characteristics. It is inferred that decohesion of precipitate/ matrix interfaces in the matrix results in the formation of quasi-cleavage fracture of the H900 specimen tested in the H2 S solution.
4. Conclusions 1. Regardless of testing environments, the peak-aged (H900) specimen demonstrates the highest NTS among all specimens. All specimens display low NTS loss in gaseous hydrogen but significantly suffer from sulfide SCC, especially for the SA specimen. 2. Hydrogen permeation properties of various specimens are strongly related to the microstructure of 17-4 PH stainless steel. Both high hydrogen diffusivity and low hydrogen solubility of the SA specimen provide the mechanism of rapid transportation of hydrogen atoms to the highly stressed region with little interference. Therefore, severely deteriorated NTS is observed for the SA specimen tested in the H2 S solution. 3. Microstructural observation of the over-aged (H1025) specimen reveals those blocky austenites together with Cu-rich precipitates uniformly distributed in grain interior. Dense and coarse precipitates are delineated along prior austenite grain boundaries. The H1025 specimen with the lowest Deff and highest Capp among all specimens can be attributed to the presence of numerous trapping sites prohibiting the inward diffusion of hydrogen atoms. Grain interior has higher resistance to
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hydrogen embrittlement than grain boundaries. Hydrogen atoms are trapped along grain boundaries and lead to the formation of intergranular fracture for H1025 specimen tested in the H2 S solution. 4. Fine and homogeneously distributed precipitates in the H900 specimen account for its superior NTS tested in the H2 S solution. Hydrogen atoms are trapped predominantly in the precipitate/matrix interfaces. The decohesion of these interfaces results in the formation of quasi-cleavage fracture of the H900 specimen tested in H2 S solution. 5. Because the microstructure in the SA condition is very susceptible to hydrogen embrittlement in the H2 S solution, the application of laser-surface-annealing treatment on the peak-aged specimen (LA specimen) cannot improve its resistance to sulfide SCC.
Acknowledgements The authors gratefully acknowledge the financial support from the National Science Council of ROC (contract no. NSC 89-2216-E-019-001).
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