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Figure from:Stress Corrosion Cracking of Pipelines

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‘igure 2-1 Three essential factors required for the occurrence of SCC.

Source: Craig and Lane [2005]. TABLE 2-2 Types of Metals That Are Susceptible to Hydrogen Damage

Figure 2-4 Typical S—N curve for fatigue and corrosion fatigue measurements of a material. Corrosion fatigue (CF) is a conjoint action of a cyclic stress and a corrosive environ- ment that results in material failure. The basic role of the corrosive environment is to decrease the life of the component. The fatigue and corrosion fatigue strength of a material is often determined by the S—N curve, that is, the stress (S)-number of stress cycles to failure (NV) curve, which is a stress against the logarithm of the number of stress cycles to a failure curve, as shown in Fig. 2-4. When a corrosive environment is present, the ability of the material to resist applied stress a particular number of cycles of stress is decreased; that is, the stress required for both crack initiation and propagation is lower in corrosive environments. The crack growth rate can be much higher in a corrosive environment than it is in a noncorrosive environment. Therefore,

eee ee ges Oe ge es ae ee Se ee At the end of stage I, the crack propagates macroscopically at 90° to the applied stress axis in stage II of the process. Multiple slip processes occur, leading to blunting of the crack tip and the formation of striations on the fracture surface. The striations we perpendicular to the crack growth direction and look like waves on the ocean surface. The striations represent one burst of crack growth but may take more than me cycle to form. Striations are characteristic of the fatigue process. The time for rack growth in stage II may take from 75 to 50% of the fatigue life. TTT 1

Figure 2-6 Effect of corrosion on fatigue crack propagation in stage II.

Source: Craig and Lane [2005]. TABLE 2-3 Specific Features of SCC, CF, and HIC

Figure 3-1 Chronology of SCC events in Canadian pipelines up to 1996. (From National Energy Board [1996a].)

Figure 3-2 Intergranular SCC in pipelines under high-pH conditions (magnification: 250 times). (From National Energy Board [1996a].)

tend to coalesce or link together to form longer cracks. At some point these cracks may reach a critical depth and length combination that can result in a rupture. Stress corrosion cracks do not need to fully penetrate a pipe wall for a rupture to occur (i.e., a shallow crack may reach a length that becomes critical). The strength and ductility of the remaining wall determines the critical size at which crack behavior changes from a slowly growing stress corrosion mechanism to an extremely rapid brittleness. A leak will occur if a crack grows through the pipe wall before it reaches a critical length for rupture. Teale pan laceanreAe A PRemMman mhanarmanmrn in mnmaliina CET Pearl -wnoleceanre:

Figure 3-4 Schematic diagram of the life cycle of pipeline SCC. (From National Energy Board [1996a].) Furthermore, the cracks in colonies can be closer together or farther apart, which makes the crack spacing either sparse or dense. Circumferential spacing equal to 20% of the wall thickness between cracks is used as the criterion that distinguishes between sparse and dense. Generally, cracks spaced circumferentially closer than 20% of the wall thickness tend to go dormant, whereas cracks spaced farther apart (at distances greater than 20% of the wall thickness) can continue to grow [Leis, 1995]. Thus, crack coalescence poses a potential mechanism affecting crack growth. The most obvious differences between the two forms of SCC include the tem-

TABLE 3-1 Comparison of Characteristics of Two Types of Pipeline SCC that supports nearly neutral-pH SCC may contain diluted bicarbonate ions and other species, such as chloride and sulfate ions, while high-pH SCC commonly occurs in environments that contain concentrated carbonate and bicarbonate ions [National Energy Board, 1996a]. These environments result from the combined effect of coating performance failure, CP penetration level, and soil conditions. However, when coating is damaged or missing, a pipeline is exposed to soil directly. The soil environment may result in pipeline SCC.

SHUTTLE HE SAUVOLY (Gd kK SICY ditd dys Hldai KEG Ullll ). During a wet—dry cycle, the thickness of the solution layer trapped under a dis- bonded coating decreases, due to the evaporation of water. Reduction in the solution thickness facilitates oxygen diffusion and reduction on the steel surface. Simulta- neously, the solution concentration increases with the thinning solution thickness, resulting in a reduction in oxygen solubility. The former enhances steel corrosion by increasing the cathodic partial reaction, resulting in a negative shift in the corrosion potential of the steel, while the latter probably causes a positive shift of corro- sion potential due to the increasing oxidizing ability of a concentrated carbonate

Figure 3-6 Time dependence of stress and corrosion potential of 3.918% prestrained X100 steel during tensile testing at a strain rate of 1x 10~¢s~' in a nearly neutral-pH solution. (From Xu and Cheng [2012].)

surface. A small anode/big cathode geometry is thus generated to accelerate localized anodic dissolution of the steel, and the electrons generated are consumed by the cathodic water reduction. Consequently, the cathodic reaction is enhanced to maintain the charge neutrality. Therefore, the porous deposit layer enhances hydrogen evolution and thus hydrogen-induced SCC susceptibility.

A buried pipeline is subjected to stress of various types from various sources, all of which contribute to SCC. The operating pressure of pipelines is normally the greatest source of stress—hoop stress—on the pipe wall. The soil that surrounds the pipe can move and generate longitudinal stress. Pipe manufacturing processes, such as welding, can create residual stresses.

Figure 3-10 Distribution of stress on a steel pipe with a corrosion effect (top: no soil strain, middle: tensile soil strain, bottom: compressive soil strain. From left to right: operating pressures of 15, 20, and 25 MPa). (From Xu and Cheng [2012].)

en i eT Te eee eT RRR SE SERENA ene) Maen) Mee ne Se oo ee It is also expected that there exist effects of the rate of pressure change or strain rate n crack initiation and growth. Generally, the slower frequencies give faster per-cycle rowth rates [Canadian Energy Pipeline Association, 1996b]. At low frequencies, rack growth per cycle is higher since there is more time for the environment to nteract with the crack, increasing the growth rate. Here, a local additional potential LAP) model is introduced to illustrate the critical role of strain rate in the SCC of ipeline steels [Liu et al., 2009b]. ii ee ee aoa le is

To establish the relationship between the LAP and the strain rate, the number of slip dislocations is derived. Figure 3-13 is a simplified view of a crystalline grain on a steel surface, where the upper surface of the grain is assumed to be exposed to the solution. The solid rectangular lines define the initial boundary of the grain, where L and d are the initial size and depth of the grain, respectively. The dashed lines constitute a new grain boundary after the deformation of AL and Ad. A two- dimensional diagram was used for simplification. W is the thickness of the grain, and its deformation is AW. It is assumed that the time period for L changing to L+ AL is t, the average depth of the source dislocation is Haye, the dislocation density is p, the generating frequency of new dislocations is f, and the shear stress is the only driving force to activate dislocations. Moreover, vy, is assumed to be the minimum mobility velocity of dislocations on the steel surface, and dislocation with a mobility velocity higher than v;, will reach the steel surface within time f. The strain. €. is defined as

Apparently, both the density, mo, and mobility, uvy,, of dislocations on the steel surface (i.e., dislocation emergence points) depend linearly on the strain rate, as demonstrated in Eqs. (3-10) and (3-11), respectively. iad ia i dain) ain 1 emia: iia” Jae Figure 3-14 shows the normalized dependence of the mobility and density lislocation emergence points on the logarithm of strain rate. It is seen that as th strain rate increases, the density of the dislocation emergence points increases, whic vould enhance the local electrochemical activity of the steel electrode due to the LA -ffect. Simultaneously, the mobility of the dislocation emergence point increases wit strain rate. When the strain rate is sufficiently high to result in a sufficiently hig nobility of dislocation emergence points on the steel surface, the reactive specie ransporting from the solution do not have an opportunity to adsorb at these activ sites for an electrochemical reaction. Therefore, there is a dual role for strain rate i electrochemical activity of the steel under straining. With respect to the effect of th strain rate on SCC susceptibility, it is expected that a maximum value phenomeno vill be observed. Detailed experimental results about the effect of strain rate on th susceptibility of pipeline steels to SCC are included in Chapter 6. —=* Tt °4 r aa an 1

after 32 days of immersion, as shown in Fig. 4-la. The cathodic current densities obtained at a potentiostatic polarization of —1.5 V(SCE) shifts negatively with time (Fig. 4-1b). The Nyquist diagrams show regular semicircles, the size decreasing with time, indicating increasing hydrogen evolution at the cathodic polarization applied. Apparently, CP could permeate through FBE, as indicated by the electrochemical responses measured. ENE LE LE Furthermore, newly developed multiple-layer coatings, such as high-performance composite coating (HPCC) is not permeable to CP, as demonstrated in Fig. 4-3, where the corrosion potential, potentiostatic polarization current, and Nyquist diagrams are measured on an X70 steel electrode that is separated from RE and CE by the HPCC membrane in two cells. It is seen that although the corrosion potential of the steel decreased with time (Fig. 4-3a), it never reached the corrosion potential of bare steel [i.e., —0.87 V(SCE)], in the soil solution. Even at a potentiostatic polarization of —1.5 V(SCE), the anodic current density is obtained, and there is little effect of time on the current density measured (Fig. 4-3b). The Nyquist diagrams measured at various times are highlighted by the large semicircle. Due to the random data obtained in the low-frequency range, the semicircles are not complete. Moreover, there is little change in semicircle size with time, and the impedance behavior of HPCC shows a purely capacitive behavior. According to coating performance evaluation criteria developed previously [King et al., 2002], HPCC, behaving as an ideal capacitor, is capable of providing full protection of the steel. Apparently, HPCC is impermeable to the CP applied, due to its extremely high impedance. In addition to the intrinsic property of coatings, the coating application practice

VUITOHL WOUIG TCTCds© WG SUSCOPUDHILY OF COaAUNE LO CaAulOQIe GISUONUGMICHL, The CP applied on a coating holiday is also partially shielded to reach the base of he holiday, due to the geometrical restraint, such as one with a diameter as small as ).2 mm [Dong et al., 2008]. Figure 4-5 shows Nyquist diagrams measured by LEIS with a microprobe positioned at the 0.2-mm defect on an HPCC-coated X70 steel specimen under various cathodic potentials in a carbonate—bicarbonate solution. For a -omparison, the EIS diagrams were also measured on bare X70 steel which was under -athodic potentials identical to those in Fig. 4-5, and the results are shown in Fig. 4-6.

Figure 4-5 Nyquist diagrams measured by LEIS with a microprobe positioned at the defect (0.2 mm in diameter) on HPCC-coated X65 steel under various cathodic potentials. (From Dong et al. [2008].)

LPessler Cl dl., LAs]. As summarized, it seems that the soil chemistry is not related directly to the occurrence of pipeline SCC. Instead, an electrolyte trapped under disbonded coating results in pipeline SCC. However, tests have shown that the high susceptibility of pipeline steels to SCC in soil-extracted solutions is always associated with a high hydrogen permeation current in the solution [Cheng, 2007a]. This indicates clearly that the soil chemistry would contribute to the occurrence of pipeline SCC and even change its mechanism either by hydrogen evolution and permeation processes or by other potential processes, such as the microorganism culture.

n organogenic formation of calcium carbonate or as a product of microorganism ctivity and their remains [Ott, 1998]. Moreover, all cases of SCC of buried pipelines nvestigated showed that soil microorganism activity, including SRB, heterotrophic, nd acid-forming bacteria colonies, was 10 times higher than that of soil in an average tate [National Energy Board, 1996].

Figure 4-10 Von Mises stress and strain distribution along an X100 steel specimen that is under a load of 200 N. (From Xu and Cheng [2012b].) Hoop stress resulting from internal pressure can be controlled appropriately. How- ever, residual stresses from manufacturing and welding, surface loading from heavily loaded trucks, longitudinal stress from ground movement, and stress concentration at dents or gouges or corrosion pits can enhance the stress level exerted on pipelines remarkably, but they are beyond the control of pipeline operators. However, for deformation-induced stress or other residual stresses that are elastic in nature, there is no effect on electrochemical corrosion of steel [Xu and Cheng, 2012b]. The critical failure strain of corrosion scale formed on a steel surface is calculated to be 0.00357 to 0.00417, which is higher than the maximum strain of 0.0029 under an applied stress of 582 MPa on the specimen, as shown in Fig. 4-10. Thus, the corrosion scale formed does not fracture during elastic loading. There is no difference in the steady-state corrosion potential detected upon application of an elastic tensile or compressive

orrosion would be significant [Xu and Cheng, 2012c]. Since hydrogen is involved in nearly neutral—-pH SCC, any factors affecting hydro- en permeation and absorption would affect the SCC process. Tests showed that an pplied stress that is equivalent to the gross hoop stress in a pipe wall under an oper- ting internal pressure of 70 bars significantly accelerates hydrogen absorption to <52, X70, and X100 steels [Capelle et al., 2010]. The difference between subsurface \\ydrogen concentrations in unloaded and stressed steel can exceed several-fold. Thus, tressed pipelines would be able to self-accelerate hydrogen permeation, contributing o the occurrence of cracking.

where ip is the current density at defect, AE the LAP at defect, 8, the cathodic Tafel slope, and iy the current density at a nondefect area (i.e., an intact area). If cathodic reaction is controlled by a mass-transport step, cathodic reactants such as H*, HCO3~, and H2CO; will be consumed more rapidly at defects than in an intact area. Figure 4-15a is a schematic view of the double-charge layer on a steel surface under a constant cathodic potential field, where both the defect and nondefect areas are associated with a uniform double-charge layer. All cathodic reactions are in equilibrium with the mass-transport process of reactants, and the steel is under full CP and does not corrode.

Figure 4-16 Polarization curves measured on an X70 steel electrode at various potential sweep rates in a nearly neutral—pH soil solution. (From Liu et al. [2011].)

Figure 4-18 Cathodic Tafel slope as a function of the upper switch potential during CV mea- surements in Fig. 4-17. (From Cheng and Niu [2007].) and 4-18 show cyclic voltammograms measured on X70 pipeline steel with various upper switching potentials at a potential sweep rate of 50 mV/s in 0.01 M NaHCO; solution, and the cathodic Tafel slope as a function of the upper switch potential, respectively. It is seen that the cathodic Tafel slope becomes less negative with an increase in upper switch potential; that is, the absolute value of the cathodic Tafel slope decreases with increasing upper potential limit. The cathodic Tafel slope fot hydrogen evolution reaction, b,, is related to the charge-transfer coefficient, a, by

where By is the charge-transfer coefficient of the charged steel. The anodic Tafel slopes, b,, can be determined by fitting the measured anodic polarization curves on steel with various hydrogen-charging current densities, and the results are shown in Fig. 4-19. It is seen that the values of b, are approximately independent of the charging current density. It is thus reasonable to assume that 8 is approximately equal to By. Reorganize Eqs. (4-13) to (4-15):

Figure 4-21 Polarization curves of X70 steel measured in a thin NS4 solution layer after immersions of 1 and 24 h: (a) 60 wm; (b) 90 wm; (c) 140 wm. (From Fu et al. [2009].)

where AU is the internal energy change; AS is the entropy change; W,,, is the mola weight; 01, 02, and o3 are the principal stresses; p is the density, E is Young’: modulus, o, is the volume stress, and Vy is the average volume of hydrogen in steel In particular, ky is the effect of hydrogen on the anodic dissolution rate in the absenc: of stress (i.e., the free corrosion rate of Fe in a deoxygenated, nearly neutral—pk solution), k, is the effect of stress on anodic dissolution in the absence of hydroget (i.e., a pure anodic dissolution-based cracking mechanism), and ky, is the synergisti effect of hydrogen and stress on the anodic dissolution at the crack tip.

Figure 4-22. SVET measurement on a precracked X70 steel specimen in an NS4 solution under a tensile force of 3000 N. (From Tang and Cheng [2009].) Furthermore, the hydrogen effect factor at a crack tip of unstressed steel at a charging current density of 20 mA/cm? is 1.3 [Zhang and Cheng, 2009b]. Simi- larly, the SVET current density maps are measured on a cracked steel specimen that is under loading and various hydrogen-charging current densities, as shown in Fig. 4-23. The synergistic effect factor, determined based on Eq. (4-27) and an extrap- olation of the charge-transfer resistance at the crack tip to a charging current density of 20 mA/cm?, is assessed to be approximately 4.0 [Tang and Cheng, 2011]. The total effect of hydrogen and stress on anodic dissolution of a precracked steel specimen under nearly neutral—pH conditions is determined as

han the 20-fold difference compared to the dissolution-based SCC propagation rate letermined by electrochemical scratch electrode testing. Therefore, in the absence of racks or cracklike flaws, the stressing and hydrogen-charging conditions employed re incapable of incurring sufficient anodic dissolution on steel in a nearly neutral-pH nvironment. However, when cracks are present, the total effect factor determined is

Figure 5-4 Schematic diagram of a U-bend X70 steel specimen and a LEIS scanning area. (From Li and Cheng [2008].)

Figure 5-6 Optical image of the morphology of a U-bend X70 steel specimen after testing in a carbonate-bicarbonate solution. (From Li and Cheng [2008].)

Figure 5-5 LEIS plots measured at three points across the various stressed areas marked in Fig. 5-4, where x = 0, 1.8, and 3.6 mm, respectively, with a constant y = 1.35 mm ina concentrated carbonate—bicarbonate solution. (From Li and Cheng [2008].)

Figure 5-8 Cyclic polarization curves measured on uncharged and precharged X80 steel elec- trodes in a carbonate-bicarbonate solution containing 0.01 M NaCl. (From Xue and Cheng [2010].)

Figure 5-9 Polarization curves of X70 steel measured in a concentrated carbonate—bicarbonate solution at a thickness of 30 .m as a function of applied stress. (From Fu and Cheng [2010].)

The crack tip works as a stress raiser when tensile stress is applied, whereas the crack wall is free of significant stress concentration. Actually, there is a distinct current density difference in the presence and absence of applied stress, as indicated in Fig. 5-10. The crack tip is subjected to a high dissolution rate, due to the periodic film rupture resulting from applied stress and local stress concentration, while the crack wall may still be in a passive state. As a consequence, the dissolution of steel at the crack tip will be enhanced further. Since the solution electrochemistry and stress state at the crack tip are quite different from those in the region ahead of the crack, the passive film formed at the crack tip would be different from that formed in the surrounding region, affecting the crack propagation. Ac chown in Fio 5-11. svcetematic electrochemical characterization has heen con-

and the stability of passive film formed at a crack tip and away from the crack tip (i.e., point A in Fig. 5-11) in concentrated carbonate—bicarbonate solutions [Zhang and Cheng, 2010]. Figure 5-12 shows Mott—Schottky plots measured at a crack tip and at point A when a steel specimen is under a 1500-N force in a concentrated carbonate-bicarbonate solution. An approximately linear relationship is observed in both C7? ~ Epp (C is the capacitance of the space-charge layer of a passive film, and Epp is the applied potential) plots with a positive slope, indicating that the passive film formed on the steel surface is an n-type semiconductor. Moreover, the slope of the line obtained at the crack tip is obviously smaller than that measured at point A. Donor densities in passive films, which are inversely proportional to the Mott-Schottky slope, formed at the crack tip and point A are 5.19 x 107! and 2.68 x 10?! cm™?, respectively, indicating the highly doped structure of the passive film. The thicknesses of the space-charge layers at the crack tip and point A are 4.16 and 5.34 A, respectively. From the parameters derived, the passive film formed at the crack tip is more unstable, as indicated by the higher donor density and thinner space-charge layer than those formed at point A. T iG senses ae them: weedde eee eed serntcece oi sek Sls. wearers Sa ee “ eee re

The galvanic coupling effect also contributes to enhanced anodic dissolution at the crack tip; that is, the crack tip serves as an anode and the adjacent region as a cathode [Zhang and Cheng, 2010]. The corrosion potential is more negative at the crack tip than at point A, the region ahead of the crack. The galvanic coupling effect causes the crack tip and the adjacent region to be polarized to a galvanic potential E,. The anodic current density at the crack tip, i;, and the cathodic current density at the adjacent region from the crack tip, i2, can be written as adjacent region from the crack tip, i2, can be written as where A; and A» are the areas of the crack tip and the region ahead of the crack, respectively. For the region away from the crack, there is cathodic polarization and, thus, Ey — Econ < 0. When E, is quite different from Eeorr1 and Econ2, e £8 Bea )/Ber and e~e~Feon2)/Bu2 are negligible. Combining with Eqs. (5-16) to (5-18), J, can be determined:

When pipelines are subjected to a seasonal wet—dry cycle, CP may reach the steel surface in wet seasons, creating a high-pH carbonate-bicarbonate solution under a disbonded region [Perdomo et al., 2001; Song, 2008]. In the dry seasons, evaporation of water generates concentrated carbonate—bicarbonate solutions, and CP becomes inaccessible to the steel, due to the increased soil resistivity. Consequently, the poten- tial of pipeline steel will shift to a positive direction and may relocate in the potential range in which SCC is susceptible. Moreover, the stability of passive film formed in the absence of CP is much higher than that formed under CP. Therefore, crack growth is more likely to occur in a dry season than in a wet season.

Figure 6-3 SEM view of crack morphology (a) and EDX spectra on inclusion A (b) and inclusion B (c). (From Liu et al. [2009a].)

Conversely, inclusions enriching silicon are easy to deform, effectively relieving the residual stress [Garet et al., 1998]. Moreover, SiO2-enriching inclusions are sphere-shaped and are prone to be in a stable state. There is a relatively small local lattice deflection around these types of inclusions. Thus, there is usually no crack associated with inclusions enriching silicon, as shown in Figs. 6-3 and 6-4.

Details of the dependence of cathodic current density on the strain rate are shown on a cathodic polarization curve as a function of the logarithm of the strain rate at an individual cathodic potential in Fig. 6-7. It is shown that at an individual stress level there is a maximum of cathodic current density at a strain rate of 10~° s7!. However, when the potential is up to —1200 mV(SCB), the cathodic current density is approximately independent of strain rate.

Figure 7-2. Optical view of the microstructure of various zones of a welded X70 steel specimen: (a) base steel; (b) HAZ; (c) HAZ/weld metal; (d) weld metal. (From Zhang and Cheng [2009b].)

and cooling rate. During the welding process, heating and cooling cause expansion and contraction of the weld area. The adjacent base metal acts as a restraint for any expansion or contraction that the weld area undergoes. This leads to residual stress buildup when it cools down to room temperature.

Figure 7-6 Schematic diagram of a welded X100 pipeline steel specimen. (From Zhang and Cheng [2010].)

solubility value for atomic hydrogen formed by the dissociation of water vapor or by a hydrocarbon in the welding arc. The rate of diffusion of hydrogen in steel at or near its melting temperature is quite high. Therefore, a molten weld pool can rapidly pick up atomic hydrogen from the hot gas in the arc. Hydrogen atoms can diffuse rapidly from the weld metal into the HAZ of the base metal, where a significant number of traps exist to accumulate hydrogen.

Figure 7-9 SVET mapping of the corrosion current on a welded X70 steel specimen containing weld metal (0 to 2 mm), HAZ (2 to 8 mm), and base steel (8 to 10 mm) zones in a nearly neutral—pH solution. (From Zhang and Cheng [2009a].)

Figure 7-10 SVET current density map measured on a welded X100 steel specimen in a nearly neutral-pH solution. (From Zhang and Cheng [2010].)

Figure 7-11 SVET measurements on a welded X70 steel specimen (0 to 2 mm: weld metal; 2 to 8 mm: HAZ; 8 to 10 mm: base steel) in a nearly neutral—pH solution after 2 h of charging at 5 and 10 mA/cr’? of charging current densities. (From Zhang and Cheng [2009a].)

TABLE 8-1 Projects Executed with Line Pipes Made of X80 Steel Path project in the fall of 2002 and was utilized on a winter construction project at Godin Lake, Alberta, with 2 km of X100 pipe included in the loop [Mohitpour et al., 2001; Glover et al., 2004]. In 2006, IPSCO in Regina successfully produced and supplied helical-welded X100 steel pipes of 42 in. x 0.5 in. wall thickness over 2 km in length. IPSCO also supplied X100 steel pipes of 30 in. x 0.375 in. wall thickness for the Fort McKay project [Horsley, 2007]. ge ae cs a ey ee eS. ne oe eS ie ee ey

Figure 8-2 Optical view of the microstructure of X100 pipeline steel. (From Jin et al. [2010].)

possess higher Cu, Mn, Mo, and Ni contents. These alloying elements tend to increase the metal hardenability and shift the continuous cooling transformation (CCT) dia- gram line to the right, increasing the formation of low-temperature transformation microstructures such as bainite and martensite under the same cooling rate.

Figure 8-7 Crack initiated at titanium oxide and aluminum oxide inclusions in X100 steel. (From Dong et al. [2009].)

Figure 8-8 Transient hydrogen-permeation current density curve measured on X80 steel (0.5 mm in thickness). (From Xue and Cheng [201 1].)

Furthermore, grain boundaries present metallurgical nonuniformity in steels. O particular importance to pipeline SCC are the interganular and transgranular SCC that occur in high- and nearly neutral-pH environmental conditions, respectivel [National Energy Board, 1996]. Figure 8-12 shows anodic current density maps mea sured by a scanning vibrating electrode technique (SVET) and the correspondin; metallographic views of X70 pipeline steel in a nearly neutral—pH bicarbonate solu tion [Liu et al., 2010]. It is seen that the steel contains a polygonal ferrite matrix, an the current density at the grain boundary is always lower than that measured on | grain. The surface morphologies of the steel electrode after SVET measurement ar shown in Fig. 8-13a, where a corrosion pit is formed on the grain surface rather thai

Figure 8-13 SEM views of the surface morphology of an X70 steel electrode after SVET measurement in Fig. 8-12 (a), and after etching in a 4 wt% HNO 3-ethanol solution (b). (From Liu et al. [2010].)

Se SS SE Ses SA Le ee ET During dynamic strain aging, solute atoms are able to make a diffusive jump while in alloy is being deformed, and therefore this process is sensitive to temperature nd strain rate [Hosford, 2005]. Thus, dynamic strain aging usually occurs at high emperatures during plastic deformation, whereas static strain aging occurs at ambient emperatures after a long period of prestrain. Dynamic aging has some distinctive haracteristics, such as a high rate of work hardening and a serrated stress—strain ‘urve that is attributed to the locking and unlocking of dislocations by solute atoms luring deformation.

Figure 8-18 Stress—strain curves of (a) anew X100 steel and (b) an aged X100 steel specimen. (From Sanjuan Rivero! [2008].)

Figure 8-19 Microstructures of (a) anew X100 steel and (b) an aged X100 steel. (From Sanjuan Riverol [2008].)

ULI! Al Asis PLUS) SHUW) CLIT SUSE VUUTIILY LU OU AL VAUIUULY PUES. 11 fractographic views of the two specimens after SSRT at CP also confirm the effect of strain aging on the increasing SCC susceptibility of steel, as shown in Fig. 8-21. Therefore, strain aging is expected to affect (increase) the susceptibility of high- strength steels, such as X100 steel, to SCC, especially at the CP or more negative potentials. It is suspected that hydrogen evolution is enhanced at the cathodic potential and that hydrogen is involved in the cracking process. The essential roles of strain aging in enhanced hydrogen evolution and permeation in the steel and the resulting SCC susceptibility increase are to be investigated.

Figure 8-22 Distribution of plastic deformation of an X80 steel pipe containing a 15.28-mm- deep corrosion defect under various strain conditions (from left to right: no prestrain, 0.2% tensile strain, 0.2% compressive strain). (From Xu and Cheng [2012a].)

a, > aes Furthermore, analysis of the mechanoelectrochemical effect on corrosion of stressed steel has to consider the anodic and cathodic reactions simultaneously. As discussed, the anodic dissolution reaction would be accelerated by applied strain, which increases the electrochemical corrosion activity of the steel. The cathodic reaction (i.e., hydrogen evolution due to the reduction of water in a deoxygenated, neutral pH solution) would also be affected by external strain, which is demonstrated by the increasing cathodic current density, measured at —1 V(SCE) with prestrain applied on the steel in Fig. 8-26. It is attributed primarily to the redistribution of elec- trochemical heterogeneity and the increasing area of the cathodic process. An increase in slip steps, microcracks, and surface defects generated during plastic deformation would reduce the activation energy of hydrogen evolution [Gutman, 1998]. Dlawkcsx sewnce ccc lene nas tha Ameena Zaeiesrtenns Af aeweecclhba MV hkaae = wea] beter ee

Figure 8-27 Von Mises stress distribution of a flat plate X100 steel specimen under a tensile force of 2750 N. (From Xu and Cheng [2012b].)

TABLE 9-1 Classification of Primary Hazards and Threats Encountered on Pipelines

the safe operating pressure would be the lesser of the hydrotest pressure at failure or the maximum monitored operating pressure for 60 days. For ILI purposes, the safe operating pressure would be the lesser of the lowest calculated failure pressure of all ILI features divided by the safety factor or the maximum monitored operating pressure over 60 days.

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