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Effect of quenching and tempering process on sulfide stress cracking susceptibility in API-5CT-C110 casing steel
Materials Science & Engineering A 688 (2017) 378–387
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Effect of quenching and tempering process on sulfide stress cracking susceptibility in API-5CT-C110 casing steel
MARK
M. Liua, C.H. Wanga, Y.C. Daia, X. Lia, G.H. Caoa, A.M. Russellb, Y.H. Liuc, X.M. Dongc, ⁎ Z.H. Zhangc, a State Key Laboratory of Advanced Special Steel & Shanghai Key Laboratory of Advanced Ferrometallurgy & School of Materials Science and Engineering, Shanghai University, 149 Yanchang Road, Shanghai 200072, PR China b Division of Materials Science and Engineering, Ames Laboratory of the U.S.D.O.E., and Department of Materials Science and Engineering, Iowa State University, Ames, IA 50011-2300, USA c Tube & Pipe Department, Baosteel Research Institute, Baoshan Iron & Steel Co., Ltd, Shanghai 201900, PR China
A R T I C L E I N F O
A BS T RAC T
Keywords: Steel EBSD Grain boundaries Hydrogen embrittlement
Three quenching and tempering processes performed on API-5CT-C110 casing steel produced tempered martensite structures and similar mechanical properties but distinct sulfide stress cracking (SSC) behavior as evaluated by Double Cantilever Beam (DCB) testing. An as-quenched specimen tempered at 690 °C for two hours showed superior SSC behavior compared to another specimen tempered at 715 °C for one hour. The latter contained a larger fraction of low-angle boundaries (LABs) and higher values of kernel average misorientation (KAM) than those in the former. Moreover, one more quenching and tempering on the former specimen would produce better SSC resistance with a decrease in the fraction of LABs and the values of KAM. Since dislocations trap hydrogen more strongly than grain boundaries, the specimen with higher KAM values, as well as higher dislocation density, would trap more hydrogen atoms and lead to greater SSC.
1. Introduction Casing and tubing steel used in oil and gas fields with high pressure and aggressive corrosion environments requires both high strength and high corrosion and hydrogen embrittlement (HE) resistance. A common HE failure of steels associated with exposure to deep oil and gas wells containing wet hydrogen sulfide (H2S) is known as sulfide stress cracking (SSC). During the production of oil country tubular goods (OCTG), chemical composition and thermomechanical treatment processing are the main factors that can be optimized to obtain the most suitable microstructure to achieve high strength with superior SSC resistance. It is generally known that SSC resistance decreases with an increase in the strength of the steel [1]. Nevertheless, a conclusion was drawn by Zhao et al. [2] that the strength value was not suitable to evaluate SSC behaviors since SSC susceptibility is related to both strength and microstructure. They concluded that acicular ferrite possessed better SSC resistance than either ultrafine ferrite or the ferritic-pearlite microstructures. In addition, it is well known that bainitic or martensitic structures with low ductility possess relatively poor SSC resistance due to the ease with which cracking initiates and propagates in these microstructures [3]. Carneiro et al. [4] showed that
⁎
the refined and homogeneous quenched and tempered bainite/martensite microstructure had the best performance against hydrogen induced cracking (HIC) and SSC. In particular, the best microstructure of steels used for sour-gas service is commonly held to be hightemperature-tempered martensite [5]. When oilfield steels are exposed to a wet H2S-bearing medium, hydrogen atoms are produced as a result of the electrochemical reactions between the metal and H2S-containing environment [6]. The hydrogen atoms diffuse into the steel instead of combining to form molecules on the steel surface because of the retarding effect of the presence of H2S or hydrogen sulfide ions (HS-) on such reactions [7,8]. Then, atomic hydrogen can accumulate in traps, and they may result in the initiation of SSC or HIC by increasing the local embrittlement or hydrogen pressure theory [9]. Hydrogen traps are categorized as reversible and irreversible traps on the basis of the binding energy of traps with hydrogen atoms [10]. Grain boundaries, dislocations and microvoids that have low trap binding energies are considered to be reversible traps, which may release hydrogen continuously at ambient temperature. But for irreversible traps, such as non-metallic inclusions and carbonitrides, they are associated with high binding energies and trap hydrogen permanently at temperatures close to ambient [11,12].
Corresponding author. E-mail address: [email protected] (Z.H. Zhang).
http://dx.doi.org/10.1016/j.msea.2017.01.067 Received 26 November 2016; Received in revised form 20 January 2017; Accepted 21 January 2017 Available online 22 January 2017 0921-5093/ © 2017 Elsevier B.V. All rights reserved.
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solubility in materials. Previous research [17,33–38] showed that potential trapping sites for diffusing hydrogen atoms may include microstructural features such as grain boundaries, subgrain boundaries, dislocations, as well as impurity elements in solid solution and interfaces of secondary-phase particles. Hydrogen trapping can substantially influence hydrogen diffusion kinetics and alter its solubility and chemical potentials. Pressouyre et al. [39] proposed that a uniform distribution of strong and innocuous traps may mitigate HE by reducing the diffusion rate with which hydrogen accumulates at potential crack initiation sites. Though the effects of microstructure and crystallography on SSC behavior have been studied by the authors [40], the mechanism on the basis of hydrogen interactions has not been explored thoroughly. Previous studies rarely investigated SSC susceptibility by combining DCB testing with EBSD observation. In the present work, DCB tests performed on API-5CT-C110 casing steel specimens with different quenching and tempering processes are correlated with the specimens’ quantitative SSC resistance. Mechanical properties and microstructure are also compared and discussed. The GBCD and KAM value measured by EBSD patterns are then used to observe the detailed effect of different quenching and tempering processes on SSC behavior of API5CT-C110 casing steel. Finally, hydrogen interactions in steel are illustrated to reveal the underlying mechanism of SSC.
The reversible traps facilitate crack propagation and affect HIC susceptibility significantly [13,14]; but the effect of irreversible traps on hydrogen damage is uncertain. Titanium carbonitrides can act as beneficial hydrogen traps and delay cracking (including HIC and SSC) in H2S environments [15], while HIC cracks initiate at manganese sulfide inclusions by the recombination of hydrogen to gaseous molecules [16]. An HIC crack is strictly internal, does not require any external stress and develops mainly parallel to the steel surface and sometimes in a stepwise pattern. In contrast an SSC crack occurs in places subjected to applied stress from the surface of the steel and develops almost perpendicularly to the stress direction. Similar to HIC, SSC is assisted by HE; the crack is influenced by hydrogen traps. The specific role of hydrogen traps, especially reversible traps, on SSC susceptibility needs further study. Hydrogen interactions with metal have been investigated experimentally by numerous investigators using permeation measurements to provide useful information about hydrogen trapping behavior [17– 21], but the test for evaluating SSC, one form of cracking assisted by HE, is performed according to the NACE TM0177 standard. The Double Cantilever Beam (DCB) test (method D) developed by NACE (TM0177, 2005) is widely utilized to evaluate the SSC resistance of OCTG [22]. This method permits measurement of a steel's resistance to propagation of SSC and expresses it in terms of a critical stress intensity factor (KISSC). Perez et al. [23] studied the microstructure and corresponding KISSC value of a series of specimens performed with different quenched and tempered heat treatments, and found that a microstructure mixed with tempered laths martensite and a phase composed of ferrite and carbide together with fine distribution of Nb(C, N) possessed excellent resistance to SSC. However, they did not provide further, in-depth interpretation when tempered martensite produced by different quenched and tempered heat treatments displayed different SSC behaviors. Crystallographic features are also considered to play a key role in hydrogen interactions with steel. The grain boundary character distribution (GBCD) and strain located in the grains and/or interfacial regions are important to analyze the initiation and growth of HE cracking. Electron back scattered diffraction (EBSD) provides accurate characterization of grain boundaries and their local strain distributions by evaluating the misorientation angle between neighboring points in a grain [24,25]. Dhondt et al. [26] focused on intergranular stress corrosion cracking (SCC) in AA2050 aluminum alloy and concluded that the grain boundary (including low-angle boundaries (LAB), special coincident site lattice boundaries (CSL), and high-angle boundaries (HAB)), do not significantly affect crack initiation. One necessary condition for SCC occurrence is considered to be the high strain localization gradient [27]. But Arafin et al. [28] investigated API X65 pipeline by EBSD and proposed that the grain boundary character (GBC) plays a key role in intergranular SCC. They found that LAB and special CSL boundaries (Σ11, Σ13b, and probably Σ5) are crackresistant, while the random HABs are susceptible to cracking. Venegas et al. [29] also claimed that the paths for intergranular HIC propagation in API-5L-X46 pipeline steel are mostly HABs. In addition, the kernel average misorientation (KAM) is a parameter calculated from the EBSD analysis to assess strain distribution as well as dislocation density [30]. Strain concentration caused higher hardening by accumulating dislocations at grain boundaries, which could be an important effect in intergranular SCC [31]. The dislocations within the virgin martensite laths are considerably rearranged during tempering, leading to a marked reduction in the dislocation density. Long term tempering of the steel is also characterized by a decrease of dislocation density [32]. Therefore, the changes of the grain boundaries and dislocation defects during the tempering process could be analyzed by EBSD to provide useful insights into the effects of defects on SSC performance. A better understanding of material failure caused by HE cracking is desired to know the effects of hydrogen diffusion, trapping, and
2. Experimental procedure 2.1. Materials processing, tensile and SSC testing The investigated samples were cut from the same API-5CT-C110 casing steel provided by Baoshan Iron & Steel Co., Ltd., which contain 0.24 C, 0.22 Si, 0.47 Mn, 0.005P, 0.001S, 0.5 Cr, 0.72 Mo, 0.1V, 0.002 B, 0.02 Ti and 0.004 N (wt%). In order to obtain a high-temperaturetempered martensite, which is a microstructure that has a desirable combination of strength and SSC resistance, the following quenchingand-tempering processes were performed on the investigated steel. All of the experimental specimens were austenitized at 900 °C for 50 min and water quenched directly. Then, they were respectively tempered at 715 °C for one hour and tempered at 690 °C for two hours, which separately labeled as QT1 and QT2. Choosing a part of QT2 samples, marked as DQT2, reheated and then water quenched and tempered again following the heat treatment details of QT2 samples. All specimens were air cooled to room temperature after tempering. The applied heat treatment routes of the experimental specimens are illustrated in Fig. 1. The test coupons were machined to a gauge section of the standard tensile test sample for evaluating tensile properties and to ensure the
Fig. 1. Schematic drawings of the applied heat treatment routes of the experimental specimens.
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Fig. 2. Geometries of DCB test specimen.
The specimen surfaces were polished using 0.04 µm colloidal silica slurry for 4 h in a Vibratory Polisher (VibroMet 2, Buehler), then they were prepared for EBSD observation. The grain boundary and local strain distribution were investigated using a SU70 Hitachi field emission SEM equipped with an EBSD detector. The areas of analysis were scanned with a step size of 50 nm at an accelerating voltage of 15 kV and a sample tilt angle of 70°. The raw data processing was carried out using TSL-OIM-Analysis software. EBSD analysis parameters were conducted using inverse pole figure (IPF) maps, grain boundary maps, and KAM maps. KAM is a measure that quantifies the mean misorientation gradient around an EBSD point with respect to a defined set of neighbor points. For a given point the average misorientation of that point with all of its neighbors was calculated with a criterion that misorientations above 5 ° were excluded from the calculation.
testing directions were parallel to the rolling direction. A test consists of a set of three samples, and yield strength was defined to be the tensile stress required to produce 0.2% elongation under load [41]. The average values from the three specimens were taken to be the tensile properties of the sample. Susceptibility to SSC of the three candidate specimens was evaluated by using several wedge-loaded DCB samples. The dimensions of standard samples in DCB testing are illustrated in Fig. 2. A double taper wedge with suitable thickness was selected to provide armdisplacement and loading by inserting into DCB samples [22]. The tests were conducted in the environment of a circulating solution consisting of 5.0 wt% NaCl and 0.5 wt% CH3COOH dissolved in distilled water saturated with 1 atm H2S. The solution in the testing was kept in a constant hydrogen ion concentration of pH=2.7 at ambient temperature. The tests were performed for two weeks. The KISSC equation in TM0177 was experimentally found to a modified version [42] given by:
KISSC =
P⋅a(2 3 + 2.38h /a )(B /Bn )1/ B⋅h1.5
3. Results
3
(1)
The DCB tests suggested that samples with a certain quenching and tempering process produced approximately equal KISSC values. SEM and EBSD observations found that all of the samples had tempered martensite. However, it was also observed that the specimens with different quenching and tempering processes had differences in their KISSC values and crystallographic characteristics that clarify the nature of the SSC mechanism in this material. Three specimens undergoing different quenching and tempering processes were chosen from the samples and discussed in this paper.
where P is the equilibrium wedge load (before or after exposure), measured from the abrupt change in slope of load-vs.-displacement curve; a is the final crack length, quantified at the exposed crack surfaces by opening the DCB test specimen mechanically after the wedge removal; h is the height of each arm; B and Bn are the specimen thickness and the web thickness, respectively. Three samples with each heat treatment process were subjected to parallel DCB experiments under the same circumstances and conditions, the KISSC values of the samples were expected to represent their SSC resistance performance.
3.1. Mechanical properties and KISSC values 2.2. Microstructure and EBSD observation The mechanical properties and KISSC values of the specimens are listed in Table 1 and illustrated in Fig. 3. To some extent, three specimens with different heat treatments displayed similar mechanical properties, especially for yield strength (~790 MPa). Tensile properties of the specimens meet the requirements of the API SPEC 5CT
The specimens for microstructure characterizing were cut from the investigated DCB samples. After grinding and polishing, the sections of three specimens were etched, and the microstructure was examined by scanning electron microscopy (SEM, S-3400N, Hitachi). 380
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of carbides in DQT2 specimen are more uniform and finer than those in QT2 specimen. The area fractions of cementites in QT1, QT2 and DQT2 specimens were about 24%, 32% and 26%, respectively.
Table 1 Mechanical properties and KISSC values of the specimens. Specimen
Yield strength (MPa)
Ultimate strength (MPa)
Reduction of area (%)
KISSC (MPa·m0.5)
QT1 QT2 DQT2
786 794 793
879 897 864
26 25 27
28.4 33.0 34.0
3.3. EBSD observation Tempered martensite microstructures were produced by different heat treatments in the three specimens, and further investigation of crystallographic features by using EBSD data was deemed necessary. EBSD orientation and IPF maps at the cross section rolling directiontransverse direction (RD-TD) planes of the three specimens are shown in Fig. 5. Grains in the QT1 and QT2 specimens do not exhibit any special orientation distribution (Fig. 5(a) and (c)), i.e., the textures of the specimens are random (Fig. 5(b) and (d)). The grains with < 101 > ‖ND (normal direction) orientation shown in the DQT2 specimen of Fig. 5(e), based on the IPF map in Fig. 5(f)), displayed a {101} texture with intensity of 2.222. Thus it is obvious that the texture is weak. To some extent, the three thermomechanical processes do not generate special orientations in the three specimens. EBSD measurements also indicate that the grain sizes of the three specimens do not differ significantly. In this study, the GBC was distinguished by three categories: LABs up to 15 ° misorientation, CSLs relations up to ∑33, and non-CSL random HABs. Fig. 6 shows the GBC collected from EBSD maps at the surfaces of the three specimens. Fig. 7 shows the proportion of LABs, CSLs, and HABs in the regions studied in the specimens. The data suggest that specimen QT1 generated a higher number fraction (41%) of LABs than those presented in specimens QT2 and DQT2, and lower percentage (29%) of HABs existed simultaneously in specimen QT1 than those in specimens QT2 and DQT2. Moreover, specimens QT2 and DQT2 contain LABs in approximately equal proportions (~28%), and they also possess similar fractions of HABs (~35%). It is noteworthy that the number fraction of LABs was decreased with increasing KISSC values of the specimens, whereas a contrary tendency was observed for HABs. Meanwhile, an uncertain relationship between the fraction of CSLs and KISSC values was displayed in the three specimens. Fig. 8 shows the color-coded KAM distribution maps of the examined areas in the three samples. Heterogeneous KAM distributions were presented in the maps. The strain accumulated mainly at grain boundaries with higher KAM values. Furthermore, it is clear that specimen QT1 exhibited higher KAM than those in the specimens QT2 and DQT2. Fig. 9 illustrates the average KAM distribution in the studied areas versus the KISSC value of the three specimens with different quenching and tempering processes. As can be seen, a decreasing trend of KISSC value was evident with the increasing average KAM distribution.
Fig. 3. Mechanical properties related to KISSC value of the specimens.
specification [41]. Specimen QT1 tempered at 715 °C for one hour had a lower ultimate strength and relatively higher reduction of area than specimen QT2 tempered at 690 °C for two hours. Comparing to specimen QT2, a decrease in strength accompany with an increase in reduction of area occurred in specimen DQT2. In addition, the KISSC value displayed in specimen QT1 (28.4 MPa·m0.5) is lower than those of specimens QT2 and DQT2 (33.0 MPa·m0.5 and 34.0 MPa·m0.5). This indicates that the QT1 specimen was more susceptible to SSC, while superior SSC resistance was observed in the QT2 and DQT2 specimens. There seems no certain correlation between KISSC and strength or reduction of area. It may be that mechanical properties can not judge the material's SSC susceptibility unilaterally.
3.2. Microstructure characterization The microstructures of the three specimens characterized by SEM are shown in Fig. 4. It was observed that tempered martensites were produced in the three specimens; they consisted of recrystallized ferrite grains and spheroidized carbides (cementites). Ferrite located in specimen QT1 tempered at higher temperature (715 °C) mainly exhibited equi-axed grain shape (Fig. 4a), which replaced martensite lath boundaries. But in the QT2 and DQT2 specimens, a part of the ferrite observed in SEM micrographs (Fig. 4c and e) still displayed the martensite lath shapes rather than equi-axed shapes. The cementite particles with white color precipitated from martensite are distributed inside ferrite grains or at ferrite grain boundaries. Owing to the greater ease of diffusion at the grain boundaries, cementites preferentially nucleated in these locations and gradually spheroidized by a diffusion transformation during tempering. The spheroidization of cementites is promoted by the resulting decrease in surface energy due to the spherical particles holding relatively smaller surface areas than the precipitates presenting other shapes with the same volume [43]. The sizes of carbides in the three specimens were ranged from several tens of nanometers to hundreds of nanometers (Fig. 4b, d and f). A closer observation would reveal specimen QT1 contained more coarsened particles than QT2 and DQT2 specimens, and the distribution and size
4. Discussion 4.1. Effect of mechanical properties and microstructure on SSC susceptibility It is reported that the largest contribution of strengthening effect on martensite comes from the solid strengthening of carbon atoms, and the base strength of iron matrix together with grain size and dislocation density account for the measured strength is relatively smaller [44]. The three specimens with tempered martensite undergoing different quenching and tempering processes possessed approximately equal strengths; this may mainly result from their similar level of solid strengthening of carbon atoms after tempering. Therefore, the SSC susceptibility could not be unilaterally determined by strength in this case, though SSC threshold stresses start decreasing with an increase in yield strength from a certain critical level [45]. Moreover, the reduction of area was also unrelated to SSC susceptibility. Thus, the mechanical properties are difficult to account for the difference of the SSC behavior 381
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Fig. 4. SEM micrographs of (a) QT1, (c) QT2, and (e) DQT2 specimens; (b), (d) and (f) are higher magnification images of (a), (c), and (e).
4.2. Effect of crystallographic characteristics on SSC susceptibility
obtained in these specimens. The tempered martensites produced by quenching and tempering on the three specimens contain a relatively dislocation-free matrix with evenly distributed and spheroidized carbides, which have a good combination of strength and SSC resistance. Especially, according to Charbonnier et al. [46], additions of Mo, V, Ti and B in the experimental quenched-tempered steels would refine the cementite particle size and precipitate as very fine carbides. The increasing fine carbides not only improve the strength of the steel, but also create a very efficient high energy trapping potentiality for hydrogen retention that is beneficial for SSC resistance. However, the microstructures of tempered martensite in the three specimens make it still hard to reveal the cause of SSC behavior.
When the tempering temperature of as-quenched steel is above 650 °C, the grain boundaries and dislocation display a remarkable change during the recrystallization process. Ferrite begins to recrystallize at 600 °C and generates many new boundaries simultaneously according to Speich's result [47]. And Caron et al. [48] also confirmed that the total area fraction of grain boundary in unit volume was increased after a prolonged tempering at 700 °C in the Fe–0.2 C (wt%) alloy due to the change of dislocations into polygonized boundaries. In general, the boundaries and dislocations play a significant impact on the material deformation and fracture. Previous studies [28,49] proposed that LABs are more resistant to fracture than HABs. Nevertheless, the specimen QT1 containing a relatively larger proportion of LABs did not possess better SSC resistance. Besides, diffusible hydrogen along the random HABs may constitute a major cause of HE 382
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Fig. 5. EBSD orientation maps at the cross section (RD-TD) planes for (a) QT1, (c) QT2, and (e) DQT2 specimens; (b), (d) and (f) present the corresponding inverse pole figures of (a), (c), and (e).
[50]. The crack is more likely to propagate along HABs rather than pass through the grain boundaries, whereas it might tend to pass across LABs and continue its growth [51]. Thus, SSC cracks can advance along
the HABs or cut through the LABs. In the meantime, it is noteworthy that the high KAM value located mainly in the LABs or in the surrounding of LABs (Figs. 6 and 8). KAM 383
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Fig. 6. Grain boundary character distribution of (a) QT1, (b) QT2, and (c) DQT2 specimens.
stress they would promote SSC crack initiation and growth. Thus, the high density of dislocation is taken for promoting HE and displaying with lower KISSC value. The interpretation agrees with the correlation that KISSC value decreased with increasing average KAM which described in Fig. 9. The results indicated that the effect of high dislocation density on promoting HE is stronger than the impact of LABs inhibited fracture. 4.3. Hydrogen interactions on the mechanism of SSC The influence of grain boundaries and dislocations on SSC was further analyzed on the basis of the hydrogen's interactions with the steels. It is well known that the poisonous effects of H2S and HS- cause atomic hydrogen produced by the cathodic reaction to diffuse into steel rather than undergo a recombination reaction on the steel surface [7,8,52]. Several studies [53–56] have confirmed that hydrogen diffusion is accelerated along the grain boundaries by a mechanism of shortcircuit diffusion, which affects HE in a material. And the finer grain size and higher dislocation density in steel exhibited a significantly lower diffusivity [12]. However, HE is not only influenced by hydrogen diffusion, but also by hydrogen trapping and solubility. The classification of trap site types according to their physical nature provided by
Fig. 7. Number fraction of grain boundary character distribution in the specimens.
maps (Fig. 8) indicated the density of dislocation concentrated in the specimen QT1 is significantly higher than that stored in the specimens QT2 and DQT2. Hydrogen atoms would be trapped in the dislocations and accumulate to reach a crack tip, then in the presence of external 384
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Fig. 8. Kernel average misorientation maps of (a) QT1, (b) QT2, and (c) DQT2 specimens.
Fig. 10. Model for energy levels of the sites when hydrogen occupied. En: Diffusion activation energy in normal lattice; Es: Saddle point energy around the trapping site; Eb0 (Ed0): Binding energy of hydrogen with the grain boundary (dislocation); Eb (Ed): Activation energy needed to escape from the grain boundary (dislocation) trapping site.
Fig. 9. The average KAM distribution in the studied areas versus the KISSC value of the three specimens.
Pressouyre et al. [33] for hydrogen in iron suggested that vacancies, alloying elements, dislocations, interfaces, and microvoids could serve as possible trap sites. The energies of hydrogen evolution from different 385
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high KAM values containing high dislocation density would be more susceptible to SSC.
trap sites are unequal. The activation energy needed to escape from trapping sites is regarded as the sum of the interaction energy between trapping site and hydrogen (binding energy) and the saddle point energy, which is usually assumed to be identical to the diffusion activation energy in iron [57]. The activation energies for hydrogen evolution from grain boundaries and dislocations in pure iron are reported to be 17.2 KJ/mol and 26.8 KJ/mol [58], respectively. For the tempered martensite in the present work, Wei's experimental results [59] also confirmed that the interaction of grain boundaries with hydrogen is weaker than that between dislocations and hydrogen. Based on the aforementioned results, Fig. 10 shows the energy levels of interstitial sites, grain boundary trapping sites, and dislocation trapping sites where hydrogen resides. As shown in the diagram, the activation energy for hydrogen desorption from a trapping site is much higher than the energy for diffusion in normal lattice sites. And the energy for hydrogen evolution from trapping sites in grain boundaries is lower than that in dislocation sites. So dislocations affected HE more significantly than did grain boundaries owing to its stronger hydrogen trapping effect. In addition, the edge dislocations cores and/or geometrically necessary dislocations act as the major trapping sites and the total trapped hydrogen along these dislocations may reach up to 96% of the total hydrogen concentration [60]. Trapping increases the solubility of hydrogen and decreases the diffusivity [61]. As a consequence, the condition of specimen QT1 held at 715 °C for one hour with higher dislocation density made it easier to absorb and pile up hydrogen atoms. When the hydrogen content was beyond the critical value, hydrogen would promote dislocation motion and result in localized plastic deformation, then the mobile dislocations would enhance the transportation of hydrogen atoms to plastic zones and finally cause an increase in HE susceptibility [2,62,63]. Moreover, in view of the hydrogen-enhanced decohesion mechanism, it is expected that hydrogen accumulation at locations of high stresses result in weakening of Fe–Fe bonds beyond a critical hydrogen concentration, leading to fracture [64,65]. Similarly, the steel would be more susceptible to SSC when its dislocation defects for hydrogen atom trapping are increased. The more favorable SSC performance of the specimen DQT2 could be explained by its low dislocation density. Thus, the dislocation density of a specimen after its tempering treatment plays a major role in determining its SSC behavior.
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Albarrán, Electrochemical technique applied to evaluate the hydrogen permeability in microalloyed steels, Int. J. Hydrogen Engery 30 (2005) 1333–1338. [21] ISO 17081: 2004 (E), Method of measurement of hydrogen permeation and determination of hydrogen uptake and transport in metals by an electrochemical technique, ISO, Switzerland, 2004. [22] NACE standard TM 0177-96 Laboratory testing of metals for resistance to sulfide stress cracking in H2S environments, NACE Int., Houston, Texas, USA, 1996. [23] T.E. Perez, G. Echaniz, C. Morales, The effect of microstructure on the KISSC low alloy carbon steels, Corrosion/1998, NACE, San Diego, CA, Paper No. 98120. [24] M. Kamaya, Y. Kawamura, T. Kitamura, Three-dimensional local stress analysis on grain boundaries in polycrystalline material, Int. J. Solids Struct. 44 (2007) 3267–3277. [25] E. Demir, D. Raabe, N. Zaafarani, S. Zaefferer, Investigation of the indentation size effect through the measurement of the geometrically necessary dislocations beneath small indents of different depths using EBSD tomography, Acta Mater. 57 (2009) 559–569. [26] M. Dhondt, I. Aubert, N. Saintier, J.M. Olive, Effects of microstructure and local mechanical fields on intergranular stress corrosion cracking of a friction stir welded aluminum-copper-lithium 2050 nugget, Corros. Sci 86 (2014) 123–130. [27] T. Couvant, L. Legras, C. Pokor, F. Vaillant, Y. Brechet, J.M. Boursier, P. Moulart,
5. Conclusions In the present study, SSC susceptibility of the API-5CT-C110 casing steel performed with different quenching and tempering process were evaluated and compared by NACE DCB testing. The roles of grain boundaries and dislocations on SSC have been analyzed, and an indepth mechanism for understanding SSC was presented. The conclusions of this work can be summarized as following: (1) Different SSC susceptibility was exhibited in the three specimens with tempered martensite and similar mechanical properties; this was attributed to the effects of their different quenching and tempering processes on grain boundaries and dislocations. (2) The as-quenched specimen tempered at 715 °C for one hours contained higher proportion of LABs and higher KAM values than those in the specimen tempered at 690 °C for two hours. And the former showed inferior SSC behavior than the latter. What's more, one more quenching and tempering on the latter specimen would produce better SSC resistance and generate a relatively lower fraction of LABs and lower KAM values. (3) The dependence of SSC behavior on quenching and tempering process was mainly correlated with the synergistic effect of GBCs and dislocations in the steel. The dislocation affected SSC more markedly than did grain boundaries due to the stronger effect on trapping hydrogen at dislocation sites. As a result, a sample with 386
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Investigations on the mechanisms of PWSCC of strain hardened austenitic stainless steels, in: Proceedings of the 13th International Conference on Environmental Degradation of Materials in Nuclear Power Systems, Whistler, British Columbia, 2007. M.A. Arafin, J.A. Szpunar, A new understanding of intergranular stress corrosion cracking resistance of pipeline steel through grain boundary character and crystallographic texture studies, Corros. Sci. 51 (2009) 119–128. V. Venegas, F. Caleyo, J.L. Gonza´lez, T. Baudin, J.M. Hallen, R. Penelle, EBSD study of hydrogen-induced cracking in API-5L-X46 pipeline steel, Scr. Mater. 52 (2005) 147–152. T.B. Britton, S. Birosca, M. Preuss, A.J. Wilkinson, Electron backscatter diffraction study of dislocation content of a macrozone in hot-rolled Ti-6Al-4V alloy, Scr. Mater. 62 (2010) 639–642. J. Hou, Q.J. Peng, T. Shoji, E.-H. Hana, W. Kea, Effects of cold working path on strain concentration, grain boundary microstructure and stress corrosion cracking in Alloy 600, Corros. Sci. 53 (2011) 2956–2962. J. Pešička, R. Kužel, A. Dronhofer, G. Eggeler, The evolution of dislocation density during heat treatment and creep of tempered martensite ferritic steels, Acta Mater. 1 (2003) 4847–4862. G.M. Pressouyre, A classification of hydrogen traps in steel, Metall. Mater. Trans. A 10 (1979) 1571–1573. K.T. Kim, J.K. Park, J.Y. Lee, S.H. Hwang, Effect of alloying elements on hydrogen diffusivity in α-iron, J. Mater. Sci. 16 (1981) 2590–2596. R.A. Oriani, The diffusion and trapping of hydrogen in steel, Acta Metall. 18 (1970) 147–157. L.W. Tsay, W.C. Lee, W.C. Luu, J.K. Wu, Effect of hydrogen environment on the notched tensile properties of T-250 maraging steel annealed by laser treatment, Corros. Sci. 44 (2002) 1311–1327. R. Valentini, A. Solina, S. Matera, P.D. Gregorio, Influence of titanium and carbon contents on the hydrogen trapping of microalloyed steels, Metall. Mater. Trans. A 27 (1996) 3773–3780. M. Garet, A.M. Brass, C. Haut, F. Guttierez-Solana, Hydrogen trapping on non metallic inclusions in Cr-Mo low alloy steels, Corros. Sci. 40 (1998) 1073–1086. G.M. Pressouyre, Trap theory of hydrogen embrittlement, Acta Metall. 28 (1980) 895–911. M. Liu, C.D. Yang, G.H. Cao, A.M. Russell, Y.H. Liu, X.M. Dong, Z.H. Zhang, Effect of microstructure and crystallography on sulfide stress cracking in API-5CT-C110 casing steel, Mater. Sci. Eng. A 671 (2016) 244–253. API SPEC 5CT, Specification for casing and tube, ninth ed., API, Washington (DC), July 2011. R.B. Heady, Evaluation of sulfide corrosion cracking resistance in low alloy steels, Corrosion 33 (1977) 98–107. H. Bhadeshia, R. Honeycombe, Steels: Microstructure and Properties, third ed., Elsevier Ltd., United States, 2006. B. Hutchinson, J. Hagström, O. Karlsson, D. Lindell, M. Tornberg, F. Lindberg, M. Thuvander, Microstructures and hardness of as-quenched martensites (0.10.5% C), Acta Mater. 59 (2011) 5845–5858.
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Effect of quenching and tempering process on sulfide stress cracking susceptibility in API-5CT-C110 casing steel
MARK
M. Liua, C.H. Wanga, Y.C. Daia, X. Lia, G.H. Caoa, A.M. Russellb, Y.H. Liuc, X.M. Dongc, ⁎ Z.H. Zhangc, a State Key Laboratory of Advanced Special Steel & Shanghai Key Laboratory of Advanced Ferrometallurgy & School of Materials Science and Engineering, Shanghai University, 149 Yanchang Road, Shanghai 200072, PR China b Division of Materials Science and Engineering, Ames Laboratory of the U.S.D.O.E., and Department of Materials Science and Engineering, Iowa State University, Ames, IA 50011-2300, USA c Tube & Pipe Department, Baosteel Research Institute, Baoshan Iron & Steel Co., Ltd, Shanghai 201900, PR China
A R T I C L E I N F O
A BS T RAC T
Keywords: Steel EBSD Grain boundaries Hydrogen embrittlement
Three quenching and tempering processes performed on API-5CT-C110 casing steel produced tempered martensite structures and similar mechanical properties but distinct sulfide stress cracking (SSC) behavior as evaluated by Double Cantilever Beam (DCB) testing. An as-quenched specimen tempered at 690 °C for two hours showed superior SSC behavior compared to another specimen tempered at 715 °C for one hour. The latter contained a larger fraction of low-angle boundaries (LABs) and higher values of kernel average misorientation (KAM) than those in the former. Moreover, one more quenching and tempering on the former specimen would produce better SSC resistance with a decrease in the fraction of LABs and the values of KAM. Since dislocations trap hydrogen more strongly than grain boundaries, the specimen with higher KAM values, as well as higher dislocation density, would trap more hydrogen atoms and lead to greater SSC.
1. Introduction Casing and tubing steel used in oil and gas fields with high pressure and aggressive corrosion environments requires both high strength and high corrosion and hydrogen embrittlement (HE) resistance. A common HE failure of steels associated with exposure to deep oil and gas wells containing wet hydrogen sulfide (H2S) is known as sulfide stress cracking (SSC). During the production of oil country tubular goods (OCTG), chemical composition and thermomechanical treatment processing are the main factors that can be optimized to obtain the most suitable microstructure to achieve high strength with superior SSC resistance. It is generally known that SSC resistance decreases with an increase in the strength of the steel [1]. Nevertheless, a conclusion was drawn by Zhao et al. [2] that the strength value was not suitable to evaluate SSC behaviors since SSC susceptibility is related to both strength and microstructure. They concluded that acicular ferrite possessed better SSC resistance than either ultrafine ferrite or the ferritic-pearlite microstructures. In addition, it is well known that bainitic or martensitic structures with low ductility possess relatively poor SSC resistance due to the ease with which cracking initiates and propagates in these microstructures [3]. Carneiro et al. [4] showed that
⁎
the refined and homogeneous quenched and tempered bainite/martensite microstructure had the best performance against hydrogen induced cracking (HIC) and SSC. In particular, the best microstructure of steels used for sour-gas service is commonly held to be hightemperature-tempered martensite [5]. When oilfield steels are exposed to a wet H2S-bearing medium, hydrogen atoms are produced as a result of the electrochemical reactions between the metal and H2S-containing environment [6]. The hydrogen atoms diffuse into the steel instead of combining to form molecules on the steel surface because of the retarding effect of the presence of H2S or hydrogen sulfide ions (HS-) on such reactions [7,8]. Then, atomic hydrogen can accumulate in traps, and they may result in the initiation of SSC or HIC by increasing the local embrittlement or hydrogen pressure theory [9]. Hydrogen traps are categorized as reversible and irreversible traps on the basis of the binding energy of traps with hydrogen atoms [10]. Grain boundaries, dislocations and microvoids that have low trap binding energies are considered to be reversible traps, which may release hydrogen continuously at ambient temperature. But for irreversible traps, such as non-metallic inclusions and carbonitrides, they are associated with high binding energies and trap hydrogen permanently at temperatures close to ambient [11,12].
Corresponding author. E-mail address: [email protected] (Z.H. Zhang).
http://dx.doi.org/10.1016/j.msea.2017.01.067 Received 26 November 2016; Received in revised form 20 January 2017; Accepted 21 January 2017 Available online 22 January 2017 0921-5093/ © 2017 Elsevier B.V. All rights reserved.
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solubility in materials. Previous research [17,33–38] showed that potential trapping sites for diffusing hydrogen atoms may include microstructural features such as grain boundaries, subgrain boundaries, dislocations, as well as impurity elements in solid solution and interfaces of secondary-phase particles. Hydrogen trapping can substantially influence hydrogen diffusion kinetics and alter its solubility and chemical potentials. Pressouyre et al. [39] proposed that a uniform distribution of strong and innocuous traps may mitigate HE by reducing the diffusion rate with which hydrogen accumulates at potential crack initiation sites. Though the effects of microstructure and crystallography on SSC behavior have been studied by the authors [40], the mechanism on the basis of hydrogen interactions has not been explored thoroughly. Previous studies rarely investigated SSC susceptibility by combining DCB testing with EBSD observation. In the present work, DCB tests performed on API-5CT-C110 casing steel specimens with different quenching and tempering processes are correlated with the specimens’ quantitative SSC resistance. Mechanical properties and microstructure are also compared and discussed. The GBCD and KAM value measured by EBSD patterns are then used to observe the detailed effect of different quenching and tempering processes on SSC behavior of API5CT-C110 casing steel. Finally, hydrogen interactions in steel are illustrated to reveal the underlying mechanism of SSC.
The reversible traps facilitate crack propagation and affect HIC susceptibility significantly [13,14]; but the effect of irreversible traps on hydrogen damage is uncertain. Titanium carbonitrides can act as beneficial hydrogen traps and delay cracking (including HIC and SSC) in H2S environments [15], while HIC cracks initiate at manganese sulfide inclusions by the recombination of hydrogen to gaseous molecules [16]. An HIC crack is strictly internal, does not require any external stress and develops mainly parallel to the steel surface and sometimes in a stepwise pattern. In contrast an SSC crack occurs in places subjected to applied stress from the surface of the steel and develops almost perpendicularly to the stress direction. Similar to HIC, SSC is assisted by HE; the crack is influenced by hydrogen traps. The specific role of hydrogen traps, especially reversible traps, on SSC susceptibility needs further study. Hydrogen interactions with metal have been investigated experimentally by numerous investigators using permeation measurements to provide useful information about hydrogen trapping behavior [17– 21], but the test for evaluating SSC, one form of cracking assisted by HE, is performed according to the NACE TM0177 standard. The Double Cantilever Beam (DCB) test (method D) developed by NACE (TM0177, 2005) is widely utilized to evaluate the SSC resistance of OCTG [22]. This method permits measurement of a steel's resistance to propagation of SSC and expresses it in terms of a critical stress intensity factor (KISSC). Perez et al. [23] studied the microstructure and corresponding KISSC value of a series of specimens performed with different quenched and tempered heat treatments, and found that a microstructure mixed with tempered laths martensite and a phase composed of ferrite and carbide together with fine distribution of Nb(C, N) possessed excellent resistance to SSC. However, they did not provide further, in-depth interpretation when tempered martensite produced by different quenched and tempered heat treatments displayed different SSC behaviors. Crystallographic features are also considered to play a key role in hydrogen interactions with steel. The grain boundary character distribution (GBCD) and strain located in the grains and/or interfacial regions are important to analyze the initiation and growth of HE cracking. Electron back scattered diffraction (EBSD) provides accurate characterization of grain boundaries and their local strain distributions by evaluating the misorientation angle between neighboring points in a grain [24,25]. Dhondt et al. [26] focused on intergranular stress corrosion cracking (SCC) in AA2050 aluminum alloy and concluded that the grain boundary (including low-angle boundaries (LAB), special coincident site lattice boundaries (CSL), and high-angle boundaries (HAB)), do not significantly affect crack initiation. One necessary condition for SCC occurrence is considered to be the high strain localization gradient [27]. But Arafin et al. [28] investigated API X65 pipeline by EBSD and proposed that the grain boundary character (GBC) plays a key role in intergranular SCC. They found that LAB and special CSL boundaries (Σ11, Σ13b, and probably Σ5) are crackresistant, while the random HABs are susceptible to cracking. Venegas et al. [29] also claimed that the paths for intergranular HIC propagation in API-5L-X46 pipeline steel are mostly HABs. In addition, the kernel average misorientation (KAM) is a parameter calculated from the EBSD analysis to assess strain distribution as well as dislocation density [30]. Strain concentration caused higher hardening by accumulating dislocations at grain boundaries, which could be an important effect in intergranular SCC [31]. The dislocations within the virgin martensite laths are considerably rearranged during tempering, leading to a marked reduction in the dislocation density. Long term tempering of the steel is also characterized by a decrease of dislocation density [32]. Therefore, the changes of the grain boundaries and dislocation defects during the tempering process could be analyzed by EBSD to provide useful insights into the effects of defects on SSC performance. A better understanding of material failure caused by HE cracking is desired to know the effects of hydrogen diffusion, trapping, and
2. Experimental procedure 2.1. Materials processing, tensile and SSC testing The investigated samples were cut from the same API-5CT-C110 casing steel provided by Baoshan Iron & Steel Co., Ltd., which contain 0.24 C, 0.22 Si, 0.47 Mn, 0.005P, 0.001S, 0.5 Cr, 0.72 Mo, 0.1V, 0.002 B, 0.02 Ti and 0.004 N (wt%). In order to obtain a high-temperaturetempered martensite, which is a microstructure that has a desirable combination of strength and SSC resistance, the following quenchingand-tempering processes were performed on the investigated steel. All of the experimental specimens were austenitized at 900 °C for 50 min and water quenched directly. Then, they were respectively tempered at 715 °C for one hour and tempered at 690 °C for two hours, which separately labeled as QT1 and QT2. Choosing a part of QT2 samples, marked as DQT2, reheated and then water quenched and tempered again following the heat treatment details of QT2 samples. All specimens were air cooled to room temperature after tempering. The applied heat treatment routes of the experimental specimens are illustrated in Fig. 1. The test coupons were machined to a gauge section of the standard tensile test sample for evaluating tensile properties and to ensure the
Fig. 1. Schematic drawings of the applied heat treatment routes of the experimental specimens.
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Fig. 2. Geometries of DCB test specimen.
The specimen surfaces were polished using 0.04 µm colloidal silica slurry for 4 h in a Vibratory Polisher (VibroMet 2, Buehler), then they were prepared for EBSD observation. The grain boundary and local strain distribution were investigated using a SU70 Hitachi field emission SEM equipped with an EBSD detector. The areas of analysis were scanned with a step size of 50 nm at an accelerating voltage of 15 kV and a sample tilt angle of 70°. The raw data processing was carried out using TSL-OIM-Analysis software. EBSD analysis parameters were conducted using inverse pole figure (IPF) maps, grain boundary maps, and KAM maps. KAM is a measure that quantifies the mean misorientation gradient around an EBSD point with respect to a defined set of neighbor points. For a given point the average misorientation of that point with all of its neighbors was calculated with a criterion that misorientations above 5 ° were excluded from the calculation.
testing directions were parallel to the rolling direction. A test consists of a set of three samples, and yield strength was defined to be the tensile stress required to produce 0.2% elongation under load [41]. The average values from the three specimens were taken to be the tensile properties of the sample. Susceptibility to SSC of the three candidate specimens was evaluated by using several wedge-loaded DCB samples. The dimensions of standard samples in DCB testing are illustrated in Fig. 2. A double taper wedge with suitable thickness was selected to provide armdisplacement and loading by inserting into DCB samples [22]. The tests were conducted in the environment of a circulating solution consisting of 5.0 wt% NaCl and 0.5 wt% CH3COOH dissolved in distilled water saturated with 1 atm H2S. The solution in the testing was kept in a constant hydrogen ion concentration of pH=2.7 at ambient temperature. The tests were performed for two weeks. The KISSC equation in TM0177 was experimentally found to a modified version [42] given by:
KISSC =
P⋅a(2 3 + 2.38h /a )(B /Bn )1/ B⋅h1.5
3. Results
3
(1)
The DCB tests suggested that samples with a certain quenching and tempering process produced approximately equal KISSC values. SEM and EBSD observations found that all of the samples had tempered martensite. However, it was also observed that the specimens with different quenching and tempering processes had differences in their KISSC values and crystallographic characteristics that clarify the nature of the SSC mechanism in this material. Three specimens undergoing different quenching and tempering processes were chosen from the samples and discussed in this paper.
where P is the equilibrium wedge load (before or after exposure), measured from the abrupt change in slope of load-vs.-displacement curve; a is the final crack length, quantified at the exposed crack surfaces by opening the DCB test specimen mechanically after the wedge removal; h is the height of each arm; B and Bn are the specimen thickness and the web thickness, respectively. Three samples with each heat treatment process were subjected to parallel DCB experiments under the same circumstances and conditions, the KISSC values of the samples were expected to represent their SSC resistance performance.
3.1. Mechanical properties and KISSC values 2.2. Microstructure and EBSD observation The mechanical properties and KISSC values of the specimens are listed in Table 1 and illustrated in Fig. 3. To some extent, three specimens with different heat treatments displayed similar mechanical properties, especially for yield strength (~790 MPa). Tensile properties of the specimens meet the requirements of the API SPEC 5CT
The specimens for microstructure characterizing were cut from the investigated DCB samples. After grinding and polishing, the sections of three specimens were etched, and the microstructure was examined by scanning electron microscopy (SEM, S-3400N, Hitachi). 380
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of carbides in DQT2 specimen are more uniform and finer than those in QT2 specimen. The area fractions of cementites in QT1, QT2 and DQT2 specimens were about 24%, 32% and 26%, respectively.
Table 1 Mechanical properties and KISSC values of the specimens. Specimen
Yield strength (MPa)
Ultimate strength (MPa)
Reduction of area (%)
KISSC (MPa·m0.5)
QT1 QT2 DQT2
786 794 793
879 897 864
26 25 27
28.4 33.0 34.0
3.3. EBSD observation Tempered martensite microstructures were produced by different heat treatments in the three specimens, and further investigation of crystallographic features by using EBSD data was deemed necessary. EBSD orientation and IPF maps at the cross section rolling directiontransverse direction (RD-TD) planes of the three specimens are shown in Fig. 5. Grains in the QT1 and QT2 specimens do not exhibit any special orientation distribution (Fig. 5(a) and (c)), i.e., the textures of the specimens are random (Fig. 5(b) and (d)). The grains with < 101 > ‖ND (normal direction) orientation shown in the DQT2 specimen of Fig. 5(e), based on the IPF map in Fig. 5(f)), displayed a {101} texture with intensity of 2.222. Thus it is obvious that the texture is weak. To some extent, the three thermomechanical processes do not generate special orientations in the three specimens. EBSD measurements also indicate that the grain sizes of the three specimens do not differ significantly. In this study, the GBC was distinguished by three categories: LABs up to 15 ° misorientation, CSLs relations up to ∑33, and non-CSL random HABs. Fig. 6 shows the GBC collected from EBSD maps at the surfaces of the three specimens. Fig. 7 shows the proportion of LABs, CSLs, and HABs in the regions studied in the specimens. The data suggest that specimen QT1 generated a higher number fraction (41%) of LABs than those presented in specimens QT2 and DQT2, and lower percentage (29%) of HABs existed simultaneously in specimen QT1 than those in specimens QT2 and DQT2. Moreover, specimens QT2 and DQT2 contain LABs in approximately equal proportions (~28%), and they also possess similar fractions of HABs (~35%). It is noteworthy that the number fraction of LABs was decreased with increasing KISSC values of the specimens, whereas a contrary tendency was observed for HABs. Meanwhile, an uncertain relationship between the fraction of CSLs and KISSC values was displayed in the three specimens. Fig. 8 shows the color-coded KAM distribution maps of the examined areas in the three samples. Heterogeneous KAM distributions were presented in the maps. The strain accumulated mainly at grain boundaries with higher KAM values. Furthermore, it is clear that specimen QT1 exhibited higher KAM than those in the specimens QT2 and DQT2. Fig. 9 illustrates the average KAM distribution in the studied areas versus the KISSC value of the three specimens with different quenching and tempering processes. As can be seen, a decreasing trend of KISSC value was evident with the increasing average KAM distribution.
Fig. 3. Mechanical properties related to KISSC value of the specimens.
specification [41]. Specimen QT1 tempered at 715 °C for one hour had a lower ultimate strength and relatively higher reduction of area than specimen QT2 tempered at 690 °C for two hours. Comparing to specimen QT2, a decrease in strength accompany with an increase in reduction of area occurred in specimen DQT2. In addition, the KISSC value displayed in specimen QT1 (28.4 MPa·m0.5) is lower than those of specimens QT2 and DQT2 (33.0 MPa·m0.5 and 34.0 MPa·m0.5). This indicates that the QT1 specimen was more susceptible to SSC, while superior SSC resistance was observed in the QT2 and DQT2 specimens. There seems no certain correlation between KISSC and strength or reduction of area. It may be that mechanical properties can not judge the material's SSC susceptibility unilaterally.
3.2. Microstructure characterization The microstructures of the three specimens characterized by SEM are shown in Fig. 4. It was observed that tempered martensites were produced in the three specimens; they consisted of recrystallized ferrite grains and spheroidized carbides (cementites). Ferrite located in specimen QT1 tempered at higher temperature (715 °C) mainly exhibited equi-axed grain shape (Fig. 4a), which replaced martensite lath boundaries. But in the QT2 and DQT2 specimens, a part of the ferrite observed in SEM micrographs (Fig. 4c and e) still displayed the martensite lath shapes rather than equi-axed shapes. The cementite particles with white color precipitated from martensite are distributed inside ferrite grains or at ferrite grain boundaries. Owing to the greater ease of diffusion at the grain boundaries, cementites preferentially nucleated in these locations and gradually spheroidized by a diffusion transformation during tempering. The spheroidization of cementites is promoted by the resulting decrease in surface energy due to the spherical particles holding relatively smaller surface areas than the precipitates presenting other shapes with the same volume [43]. The sizes of carbides in the three specimens were ranged from several tens of nanometers to hundreds of nanometers (Fig. 4b, d and f). A closer observation would reveal specimen QT1 contained more coarsened particles than QT2 and DQT2 specimens, and the distribution and size
4. Discussion 4.1. Effect of mechanical properties and microstructure on SSC susceptibility It is reported that the largest contribution of strengthening effect on martensite comes from the solid strengthening of carbon atoms, and the base strength of iron matrix together with grain size and dislocation density account for the measured strength is relatively smaller [44]. The three specimens with tempered martensite undergoing different quenching and tempering processes possessed approximately equal strengths; this may mainly result from their similar level of solid strengthening of carbon atoms after tempering. Therefore, the SSC susceptibility could not be unilaterally determined by strength in this case, though SSC threshold stresses start decreasing with an increase in yield strength from a certain critical level [45]. Moreover, the reduction of area was also unrelated to SSC susceptibility. Thus, the mechanical properties are difficult to account for the difference of the SSC behavior 381
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Fig. 4. SEM micrographs of (a) QT1, (c) QT2, and (e) DQT2 specimens; (b), (d) and (f) are higher magnification images of (a), (c), and (e).
4.2. Effect of crystallographic characteristics on SSC susceptibility
obtained in these specimens. The tempered martensites produced by quenching and tempering on the three specimens contain a relatively dislocation-free matrix with evenly distributed and spheroidized carbides, which have a good combination of strength and SSC resistance. Especially, according to Charbonnier et al. [46], additions of Mo, V, Ti and B in the experimental quenched-tempered steels would refine the cementite particle size and precipitate as very fine carbides. The increasing fine carbides not only improve the strength of the steel, but also create a very efficient high energy trapping potentiality for hydrogen retention that is beneficial for SSC resistance. However, the microstructures of tempered martensite in the three specimens make it still hard to reveal the cause of SSC behavior.
When the tempering temperature of as-quenched steel is above 650 °C, the grain boundaries and dislocation display a remarkable change during the recrystallization process. Ferrite begins to recrystallize at 600 °C and generates many new boundaries simultaneously according to Speich's result [47]. And Caron et al. [48] also confirmed that the total area fraction of grain boundary in unit volume was increased after a prolonged tempering at 700 °C in the Fe–0.2 C (wt%) alloy due to the change of dislocations into polygonized boundaries. In general, the boundaries and dislocations play a significant impact on the material deformation and fracture. Previous studies [28,49] proposed that LABs are more resistant to fracture than HABs. Nevertheless, the specimen QT1 containing a relatively larger proportion of LABs did not possess better SSC resistance. Besides, diffusible hydrogen along the random HABs may constitute a major cause of HE 382
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Fig. 5. EBSD orientation maps at the cross section (RD-TD) planes for (a) QT1, (c) QT2, and (e) DQT2 specimens; (b), (d) and (f) present the corresponding inverse pole figures of (a), (c), and (e).
[50]. The crack is more likely to propagate along HABs rather than pass through the grain boundaries, whereas it might tend to pass across LABs and continue its growth [51]. Thus, SSC cracks can advance along
the HABs or cut through the LABs. In the meantime, it is noteworthy that the high KAM value located mainly in the LABs or in the surrounding of LABs (Figs. 6 and 8). KAM 383
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Fig. 6. Grain boundary character distribution of (a) QT1, (b) QT2, and (c) DQT2 specimens.
stress they would promote SSC crack initiation and growth. Thus, the high density of dislocation is taken for promoting HE and displaying with lower KISSC value. The interpretation agrees with the correlation that KISSC value decreased with increasing average KAM which described in Fig. 9. The results indicated that the effect of high dislocation density on promoting HE is stronger than the impact of LABs inhibited fracture. 4.3. Hydrogen interactions on the mechanism of SSC The influence of grain boundaries and dislocations on SSC was further analyzed on the basis of the hydrogen's interactions with the steels. It is well known that the poisonous effects of H2S and HS- cause atomic hydrogen produced by the cathodic reaction to diffuse into steel rather than undergo a recombination reaction on the steel surface [7,8,52]. Several studies [53–56] have confirmed that hydrogen diffusion is accelerated along the grain boundaries by a mechanism of shortcircuit diffusion, which affects HE in a material. And the finer grain size and higher dislocation density in steel exhibited a significantly lower diffusivity [12]. However, HE is not only influenced by hydrogen diffusion, but also by hydrogen trapping and solubility. The classification of trap site types according to their physical nature provided by
Fig. 7. Number fraction of grain boundary character distribution in the specimens.
maps (Fig. 8) indicated the density of dislocation concentrated in the specimen QT1 is significantly higher than that stored in the specimens QT2 and DQT2. Hydrogen atoms would be trapped in the dislocations and accumulate to reach a crack tip, then in the presence of external 384
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Fig. 8. Kernel average misorientation maps of (a) QT1, (b) QT2, and (c) DQT2 specimens.
Fig. 10. Model for energy levels of the sites when hydrogen occupied. En: Diffusion activation energy in normal lattice; Es: Saddle point energy around the trapping site; Eb0 (Ed0): Binding energy of hydrogen with the grain boundary (dislocation); Eb (Ed): Activation energy needed to escape from the grain boundary (dislocation) trapping site.
Fig. 9. The average KAM distribution in the studied areas versus the KISSC value of the three specimens.
Pressouyre et al. [33] for hydrogen in iron suggested that vacancies, alloying elements, dislocations, interfaces, and microvoids could serve as possible trap sites. The energies of hydrogen evolution from different 385
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high KAM values containing high dislocation density would be more susceptible to SSC.
trap sites are unequal. The activation energy needed to escape from trapping sites is regarded as the sum of the interaction energy between trapping site and hydrogen (binding energy) and the saddle point energy, which is usually assumed to be identical to the diffusion activation energy in iron [57]. The activation energies for hydrogen evolution from grain boundaries and dislocations in pure iron are reported to be 17.2 KJ/mol and 26.8 KJ/mol [58], respectively. For the tempered martensite in the present work, Wei's experimental results [59] also confirmed that the interaction of grain boundaries with hydrogen is weaker than that between dislocations and hydrogen. Based on the aforementioned results, Fig. 10 shows the energy levels of interstitial sites, grain boundary trapping sites, and dislocation trapping sites where hydrogen resides. As shown in the diagram, the activation energy for hydrogen desorption from a trapping site is much higher than the energy for diffusion in normal lattice sites. And the energy for hydrogen evolution from trapping sites in grain boundaries is lower than that in dislocation sites. So dislocations affected HE more significantly than did grain boundaries owing to its stronger hydrogen trapping effect. In addition, the edge dislocations cores and/or geometrically necessary dislocations act as the major trapping sites and the total trapped hydrogen along these dislocations may reach up to 96% of the total hydrogen concentration [60]. Trapping increases the solubility of hydrogen and decreases the diffusivity [61]. As a consequence, the condition of specimen QT1 held at 715 °C for one hour with higher dislocation density made it easier to absorb and pile up hydrogen atoms. When the hydrogen content was beyond the critical value, hydrogen would promote dislocation motion and result in localized plastic deformation, then the mobile dislocations would enhance the transportation of hydrogen atoms to plastic zones and finally cause an increase in HE susceptibility [2,62,63]. Moreover, in view of the hydrogen-enhanced decohesion mechanism, it is expected that hydrogen accumulation at locations of high stresses result in weakening of Fe–Fe bonds beyond a critical hydrogen concentration, leading to fracture [64,65]. Similarly, the steel would be more susceptible to SSC when its dislocation defects for hydrogen atom trapping are increased. The more favorable SSC performance of the specimen DQT2 could be explained by its low dislocation density. Thus, the dislocation density of a specimen after its tempering treatment plays a major role in determining its SSC behavior.
Acknowledgments This work was financially supported by Baoshan Iron & Steel Co., Ltd. References [1] A.W. Thompson, I.M. Bernstein, in: R.W. Staehle, M. Fontana (Eds.), Advances in Corrosion Science and Technology, vol. 7, 1979, pp. 53–175. [2] M.C. Zhao, B. Tang, Y.Y. Shan, K. Yang, Role of microstructure on sulfide stress cracking of oil and gas pipeline steels, Metall. Mater. Trans. A 34 (2003) 1089–1096. [3] B. Beidokhti, A. Dolati, A.H. Koukabi, Effects of alloying elements and microstructure on the susceptibility of the welded HSLA steel to hydrogen-induced cracking and sulfide stress cracking, Mater. Sci. Eng. A 507 (2009) 167–173. [4] R.A. Carneiro, R.C. Ratnapuli, V.F.C. Lins, The influence of chemical composition and microstructure of API linepipe steels on hydrogen induced cracking and sulfide stress corrosion cracking, Mater. Sci. Eng. A 357 (2003) 104–110. [5] M. Watkins, R. 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5. Conclusions In the present study, SSC susceptibility of the API-5CT-C110 casing steel performed with different quenching and tempering process were evaluated and compared by NACE DCB testing. The roles of grain boundaries and dislocations on SSC have been analyzed, and an indepth mechanism for understanding SSC was presented. The conclusions of this work can be summarized as following: (1) Different SSC susceptibility was exhibited in the three specimens with tempered martensite and similar mechanical properties; this was attributed to the effects of their different quenching and tempering processes on grain boundaries and dislocations. (2) The as-quenched specimen tempered at 715 °C for one hours contained higher proportion of LABs and higher KAM values than those in the specimen tempered at 690 °C for two hours. And the former showed inferior SSC behavior than the latter. What's more, one more quenching and tempering on the latter specimen would produce better SSC resistance and generate a relatively lower fraction of LABs and lower KAM values. (3) The dependence of SSC behavior on quenching and tempering process was mainly correlated with the synergistic effect of GBCs and dislocations in the steel. The dislocation affected SSC more markedly than did grain boundaries due to the stronger effect on trapping hydrogen at dislocation sites. As a result, a sample with 386
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