Further increase of ultimate tensile strength related with un-tempered ... Keywords: Hydrogen Embrittlement, HSLA steels, steel microstructure, cathodic ...
HYDROGEN EMBRITTLEMENT BEHAVIOUR OF HSLA LINEPIPE STEELS UNDER CATHODIC PROTECTION
Authors: M.Cabrini, S.Lorenzi, P.Marcassoli, T.Pastore Abstract The paper deals with Hydrogen Embrittlement (HE) under cathodic protection of steels used or recently proposed for buried pipelines of oil and gas industry. High strength low alloyed steels with different chemical composition, microstructure and ultimate tensile strength ranging between 530 to 860 MPa were considered in order to compare new steels with traditional grades. Their behaviour was assessed through constant load tests, fracture mechanic tests on modified WOL specimens, slow strain rate and interrupted slow strain rate tests in solutions simulating soil moistures and seawater, at room temperature and cathodic potentials between –2 to –0.8 V vs SCE reference electrode. All the steels were found immune under static load and HE phenomena were observed only during slow strain rate tests at potentials more negative than a critical potential, depending on steel susceptibility and strain rate. Quenched and tempered steels with high tempered martensite and very fine precipitation distribution show a better behavior than the rolled steels with banded microstructure, exhibiting the highest HE resistance. The resistance of steels characterized by banded microstructures increases with the ultimate tensile strength until about 700 MPa, than decreases. The increase of ultimate tensile strength by changing the steel microstructure, from hot rolling with coarse ferrite/pearlite to fine ferrite/pearlite microstructure, and finally very fine acicular ferrite arises the resistance of steel to HE. Further increase of ultimate tensile strength related with un-tempered martensite inside the acicular ferrite based microstructure produce a worsening of resistance. Keywords: Hydrogen Embrittlement, HSLA steels, steel microstructure, cathodic protection. Introduction High strength low alloyed (HSLA) steels are widely used in natural gas and oil transportation for buried pipelines and sealines. The pipelines are protected against external general corrosion by means of protective coatings and cathodic protection. Polarization at potentials in the range –0.8 to –1.1 V vs SCE are imposed, but very negative levels can be reached on overprotected areas close to the impressed current anodes (Shipilov, 2002, Lazzari and Pedeferri, 2006). Decades of field experience confirm the importance of cathodic protection in order to reduce the risk of perforation. However, some field ruptures pointed out the possibility of insurgence of Hydrogen Embrittlement (HE) in cathodically protected pipelines (Punter et al., 1991, Chiovelli et al., 1994, Polyakov and Kharionovsky, 1996). These phenomena were historically attributed to the presence of hard spot in heat affected zone, with altered microstructure. A great number of studies were published on this problem that was avoided with an accurate control of the steel production (Fessler et al., 1977, Carter and Hyatt, 1973). Some ruptures also pointed out the possibility to have HE in absence of hard spots, if a slow continuous plastic deformation acts on the pipeline, caused by the soil settlement or landslides (Punter et al., 1992). Similar results has been reproduced in full scale laboratory tests at cathodic potentials of –1.05 V and –2 V vs. SCE and strain rates (𝜀̇) of 10–510–6 s–1 (Cabrini et al., 1998). Furthermore, cyclic loadings due to gas pressure variation can also contribute to environmentally assisted subcritical crack growth (Shipilov, 2005, Shipilov and Le May, 2006). As a matter of fact daily pressure fluctuations in pipeline can produce cyclic loading with stress ratio (minimum/maximum load ratio) of about 0.6-0.8 at very low cyclic frequency of about 2-5 cycles per day (Gutman et al., 1977). 1
Because EAC phenomena usually increase with the mechanical properties of the steel and nowadays there is a general tendency of the Oil and Gas Industry for developing high strength grades of steel for pipeline in order to reduce wall thickness and the costs associated with pipeline projects (Tanguy et al., 2005, Seo et al., 2007, Beidoknti et al, 2009), the risk of HE could rise. Steels with tensile yield strength exceeding about 700 MPa should be verified in order to prevent cracking on structures with cathodic polarization at potentials more negative than -0.9 V vs SCE (Marchesani et al., 2005, Demofonti et al., 2008). New X100 and X120 steels are currently being considered for high pressure gas pipeline applications in the arctic (Huang et al., 2010). X100 steel has been developed for installation in projects in the northern Canada and the Japanese Sub-Sea. A number of literature works have been devoted to the study of HE under cathodic protection. The role of strain rate and potential was evidenced by Hinton and Procter (1983), Cigada et al. (1985) Saenz et al. (1986). Cigada et al. observed cracking on X65 in synthetic seawater at the usual potential for cathodic protection after high strain and with strength rate below 4∙10-5 s-1. They show the effect of strain rate on the critical potential for HE. At free corrosion potential cracking was noticed only at 10-7 strain rate. Saenz et al. (1986) found that API 5L X65 and X70 steels showed a loss of ductility during SSR tests when polarised at potential lower than -0.8 V vs SCE in 0.6 M NaCl solution. Similar results were obtained by Trasatti et al. (2005) during SSR tests on API5L X80 steel at -1 V vs SCE in substitute ocean water. The loss of ductility was due to hydrogen induced cleavage cracking which occurred as a result of the enhanced entry and transport of hydrogen in the steel produced by continuous straining. Rondelli et al. (1987) demonstrate on X65 steel under normal level of cathodic protection that also the crack propagation for HE only occur under continuous straining. The energy required for propagation is much lower than for propagation in air and decreases as cathodic polarization increases. A threshold of plastic strain for HE occurring and the role of strain rate was also confirmed by slow bending tests (Cabrini et al., 2004). The behaviour of different grades of pipeline steels was considered in previous papers (Cabrini et al., 2003, Barsanti et al., 2004, Cabrini et al., 2004). The susceptibility of the material depends on the mechanical properties, the microstructure and the chemical composition. The resistance to cracking with HE mechanism has been widely demonstrated for sour service application (Parkins, 1969). By means of thermo-mechanical controlled processing (TMCP), covering controlled rolling and accelerated cooling, steels with high strength in association with superior toughness and formability can be produced. The resulting multi-phase microstructure depends on cooling rate, ranging from polygonal ferrite-pearlite to degenerate pearlite and lath-type/bainitic ferrite (acicular ferrite). By increasing cooling rate it is possible to observe a prevalent microstructure of acicular ferrite. Accelerated cooling after controlled rolling also produces a refined final microstructure and favours the formation of low transformation temperature products, like bainite and martensite – austenite constituent (Rodrigues et al., 2000, Shanmugametal, 2007, Matrosov et al., 2008). Seamless pipelines are also produced using quenching and tempering, producing steels with tempered martensite microstructure. This paper shows the results of tests performed at the University of Bergamo on the effect of steel microstructure, cathodic potential and strain rate on hydrogen embrittlement of linepipe steels, ranging from old traditional grades to new ones recently proposed, under cathodic protection in solutions simulating soil moisture and seawater. The behaviour of 19 different linepipe steels produced either through controlled rolling and accelerated cooling or quenching and tempering heat treatments was evaluated by means of Constant Load Tests, Slow Strain Rate Tests, Interrupted Slow Strain Rate Tests and Fracture Mechanics Tests. EXPERIMENTAL Materials Table 1 summarises chemical compositions, manufacturing process, mechanical properties and microstructures of the steels. The tests were carried out on linepipe steels with Tensile Yield Strength from 430 to 770 MPa. Both SAW longitudinally welded pipes manufactured by hot rolled and controlled rolled plates and 2
seamless pipes were taking into account. The steels grades with 80/100 ksi TYS were manufactured by TMCP. Seamless pipes were quenched and tempered. The steels produced by traditional hot rolling show banded ferrite/pearlite microstructures, with polygonal ferrite and bands of coarse pearlite (Figure 1) and TYS ranging from 420 to 470 MPa. The steel manufactured by controlled rolling shows a refined ferrite grain (Figure 2) that increases yield strength and toughness with respect to the traditional hot rolled steels. The adoption of accelerated cooling after controlled rolling produces a further refinement of ferrite grain and modifies microstructure by promoting bainitic and martensitic transformations. The pearlite located between ferrite bands is progressively replaced by acicular ferrite with carbides (Figure 3), bainite and tempered martensite as cooling rate increases (Figures 4-6). The TMCP steels considered in this work show variable contents of pearlite, bainite and martensite and uneven microstructure along the thickness of pipe wall. The quenched and tempered steels have homogeneous microstructure of fully tempered martensite but different dimensions and distribution of carbides. Test solutions The tests were performed in 35g/L NaCl solution, in artificial seawater prepared according ASTM D1141-90 and in a solution proposed as the medium composition of soil moistures found under coatings of buried pipelines, named NS4 solution, usually adopted for studying Near Neutral Stress Corrosion Cracking (Delanty and O’Beime, 1992). The composition of the NS4 solution is 0.483 g/L of NaHCO3, 0.122 g/L of KCl, 0.18 g/L of CaCl2 and 0.1 g/L of MgSO4. All the tests were carried out in aerated solution at room temperature. The pH is 7 for NaCl solution, 8.3 for NS4 and artificial sea water solutions. Cathodic protection was applied through a potentiostat, in the range -0.800 ÷ -2 V vs SCE (Saturated Calomel Electrode). Specimens and tests procedure Constant load, slow strain rate tests and interrupted slow strain rate tests were performed on cylindrical tensile specimens (Figure 7) obtained along longitudinal direction, close the external surface of pipe. Before testing, the specimens were finished with abrasive paper (1200 grit) in order to eliminate the strain-hardened layer induced by machining. Finally they were polished with acetone or ether. Fracture mechanics tests were performed on 12.70 mm bolt load modified WOL specimens (Figure 8) obtained from base material with loading in longitudinal direction and propagation along transverse direction (L-T specimens), or loading in transverse direction with propagation in longitudinal direction (T-L specimens) (ASTM E399-90). Constant Load tests were performed by means of a tensile machine with a double lever arm with two–ball joints in order to achieve pure axial loading. All the specimens were loaded at the tensile yield strength estimated from loading curves of Slow Strain Rate tests in air. The SSR tests were carried out at strain rates ranging from 10–4 s–1 and 10–7 s–1 through a tensile machine, with four independent stations of loading. During SSR tests the load was monitored as a function of time in order to estimate stress vs strain curves. After the tests, the reduction of area (Z%) and total elongation (A%) were measured on specimen. Tests were performed in solution and, for comparison, in air. The occurrence of environmentally assisted cracking after tests in solution was evaluated through the maximum load max, A% and Z% normalised to the respective values measured after tests in air. Furthermore, fractographic analysis was conducted in order to detect brittle areas and secondary cracks. The Interrupted Slow Strain Rate (ISSR) tests were carried out on the same specimens of the SSR tests. During ISSR tests, the specimen was initially loaded at 105 s1 strain rate. After maximum load was reached, i.e. after 5% reduction of load from maximum value was achieved, the monotonic increase of elongation was interrupted and the specimen was left in the test solution at constant deformation for one week. At the end of exposure period, the specimen was examined at the Scanning Electron Microscope in order to detect the presence of any cracking. The fracture mechanics tests were carried out on modified wedge-opening-load (WOL) specimens using the experimental procedure described by Wei and Novak (1987). The experimental details were described in 3
Cigada et al. (1992). The specimens were fatigue pre-cracked according to requirements of ASTM E399-90 Standard, providing the maximum value of Stress Intensity Factor (Kmax) was lower than 50% of the value applied for testing (Kap). The specimens were bolt loaded and the applied load (P) was estimated through the crack opening displacement (V). The crack length was estimated by compliance measurements during last fatigue cycles of pre-cracking. At the end of tests, it was again measured by optical examination after specimen opening according ASTM E399-90 Standard. The tests were performed in the range of 73 to 88 MPa∙m1/2. Before immersion, the specimens were coated with the exception of the crack propagation zone, in order to reduce cathodic protection current. The specimens were put into a PE cell with 35 L test solutions and polarised at –1.5 V vs SCE. The tests solution was re-circulated by means of immersion pump and replaced every 15 days. The pH of the solution was monitored in order to prevent any relevant change. However, it always remained in the range 7-8. The testing time was from 5 to 18 months. RESULTS Failures were never observed during CL tests, even with cathodic polarization at very negative potential, at 1.5 and -2 V vs SCE. These results were also confirmed on precracked modified WOL specimens, which never showed any propagation from fatigue pre-crack. Also during the ISSR tests, the specimens did not show any cracks at –1.05 V vs SCE. The presence of microcracks was only observed at –1.5 V vs SCE on high strength grades produced by controlled rolling and accelerated cooling. Hydrogen embrittlement phenomena were observed in the SSR tests as a function of applied potential and strain rate. The results do not show any relevant variation in steel behaviour over the three test solutions at very negative potentials. Thus, they agree with the data of Kasahara et al. (1981) reported in the review of Shipilov (2006), that observed that the loss of steel ductility under cathodic protection was not affected by pH in the range between 4 and 12. However, the SSR tests at the potential of -1.05 V vs SCE on steel FP1 showed results in NS4 solution very similar to that in 35 g/L NaCl, but a behaviour in synthetic seawater slightly worse. Figures 9-10 show examples of stress vs strain curves obtained during SSR tests. During SSR tests, the specimens exhibited low ductility at very negative potentials and low slow strain rates, due to well-defined brittle cracks on fracture surface. As a consequence, the loading curves change but in all cases they approached to the maximum load reached during tests in air, showing an initial similar trend. According to this, secondary cracks occurred only in the necking zone. Thus, environmental embrittlement only modified the final part of the loading curve. During this phase, the final rupture of the specimens initiates from surface cracks induced by HE on heavy plastic strained material, clearly noted by SEM analysis (Figure 11), producing a decrease of ductile parameters. SEM fracture analysis was performed in order to evidence HE crack growth more than 50 µm on fracture surface and secondary cracks. On specimens showing well-defined brittle areas and secondary cracks, the Reduction of Area usually was lower than 90% of the value obtained after test in air. Fractographic examination was more effective in order to identify the insurgence of HE, as a matter of fact slight HE cracking was even remarked on specimens showing Reduction in Area close to the values measured after air tests. The Figures 12-15 illustrate the results of SSR tests. Hydrogen embrittlement occurred at potentials more negative than a critical value that depends on susceptibility of the steel and strain rate. The critical potential decreases as strain rate and resistance of steel to this phenomena increase. DISCUSSION The HE is a phenomenon caused by the penetration of atomic hydrogen inside the metal lattice. The electrochemical reduction of hydrogen ions occurs on the surface of the metal, according to the mechanism: 4
. 𝐻$ 𝑂 + 𝑒 ( → 𝐻*+, (1) . ( 𝐻*+, + 𝐻$ 𝑂 + 𝑒 → 𝐻$ (2) . . 𝐻*+, + 𝐻*+, → 𝐻$ (3) Part of the hydrogen adsorbed on the surface of the metal (Hads) diffuses in the metal and is absorbed in the lattice (Habs): . . 𝐻*+, → 𝐻*., (4) The absorbed hydrogen diffuses inside the lattice toward the zones of stress concentration and induces brittle delayed fractures after it reaches a critical concentration. In oil industry, the presence of substances that are poisons of hydrogen recombination, such as H2S, promotes high atomic concentration adsorbed on metallic surface and consequently allows high level of hydrogen diffusion in the lattice, enhancing HE cracking. For the environments considered in this research these substance are not typical. In their absence, the principal factors influencing Hydrogen Embrittlement mechanism are the cathodic polarization, the intrinsic susceptibility of the material and its mechanical and metallurgical properties, the straining conditions. The effect of the cathodic polarization is related with the role of the potential (E) on the rate of the cathodic process of hydrogen evolution from water dissociation. The adsorbed hydrogen is produced by the reaction (1). In aerated waters, at the pH of the solutions adopted for this study, this reaction is very low or negligible at the free corrosion potential of steel. The cathodic protection, necessary for the protection from the generalized corrosion, shifts the potential under the equilibrium potential for the hydrogen evolution (Eeq,H). Thus, the rate of such reaction increases according to the Tafel’s law: 67689,;
𝑖0 = 𝑖2 ∙ 10( < (5) with b that can be assumed equal to about 0.052 V at room temperature. According to this law, the rate of formation of the atomic hydrogen on metallic surface under cathodic protection increases of one order of magnitude by decreasing the potential (E) of 0.120 V. The equilibrium potential can be estimated by the first Nernst’s law, as a function of pH: 𝐸>?,0 = −59 ∙ 𝑝𝐻 (6) At the pH of soil moistures, generally from 7 to 8, or seawater, typically 8.2, the equilibrium potential is about 0.66 ÷ -0.73 V vs SCE (Figure 16). However, cathodic protection increases the pH on the protected surface enough to cause the precipitation of the carbonate scale both on marine and buried structures, for the alkalinity produced by the cathodic reaction, proportionally to the cathodic current density. In appendix, a model based on stationary diffusion flux is described for neutral solutions, without any buffer capacity, and for solutions with tendency to separate calcium carbonate. By assuming that at the hydrogen equilibrium potential the cathodic current density is equal to the oxygen limiting current, the alkalinisation is estimated to rise up to about pH 11 in well aerated solution without any calcareous deposit formation, i.e. pure NaCl solutions or solutions with low contents of calcium, magnesium and bicarbonate ions. On the contrary, the formation of carbonate deposits tends to maintain alkalinity close to the pH of saturation for carbonate precipitation, reducing the alkalinisation on the cathode. In this case, if solution contain a sufficient concentration of calcium and bicarbonate ions, the pH on the electrode is expected to increase only up to reach the necessary supersaturation for precipitation, that can be consider, in first approximation, not more than half or one unit of pH (see Appendix). On the base of this model, which is in accordance with the analysis made from other literature work since from the paper of Engell and Forchhammer (1965), it can be assumed that the pH on cathodic polarized metal approach 10.5-11 in well aerated NaCl solutions. At these pH, the thermodynamic potentials to have hydrogen evolution on the metal is in the range of -0.86 ÷ -0.89 V vs SCE. In the artificial seawater it should not arise more than 8.7-9.2 pH, giving an equilibrium potential about 1 hundred millivolts more noble. These differences might explain the slightly worse behavior observed in artificial seawater than the other two test solutions at low cathodic polarization. The hydrogen diffusion is proportional to the concentration difference between surface and metal lattice. It is mainly related with concentration on the surface, that increases as the cathodic overpotential rises. In order to reach the critical concentration for cracking, the hydrogen concentration on metallic surface should reach a 5
sufficient value to promote significant diffusion towards metal lattice. This is possible if the cathodic protection produces a sufficient polarization below hydrogen equilibrium potential. The critical concentration for cracking depends both on intrinsic susceptibility of the material and the strain conditions. As the HE susceptibility of material increases, a lower hydrogen quantity is required to have cracking and HE phenomena takes place at less negative potentials. Thus the critical potential can be assumed as an index of intrinsic susceptibility of material towards HE. The critical value of potential increases as the strain rate decreases. It lies in the range of -0.85 ÷ -1.05 V vs SCE depending on the steel microstructure and mechanical properties. Effect of plastic straining The tests carried out under constant load both on smooth tensile specimens and fatigue pre-cracked specimens did not show any initiation and propagation for HE. Therefore they show the substantial HE immunity under static loading, even for marked cathodic overprotection, at very negative potential (-1.5 V vs SCE) (Figure 16). The quantity of hydrogen is enhanced into the mechanically distorted metal lattice of the elastic strained steel but the critical hydrogen concentration in the most stressed zones was never reached, in absence of hydrogen poison (Cabrini, 1998). Different conditions occur on steels with hard spots that show cracking under constant load. Hydrogen can accumulate in the un-tempered martensitic zones, into the distorted tetragonal lattice as demonstrate by photoelectrochemical measurements on X60 steel (Razzini, 1998). In these conditions hydrogen embrittlement takes place when the specimens are loaded at constant value. During SSR tests, the continuous straining in plastic field influences the hydrogen saturation of traps into the metal. Permeation tests under slow strain rate carried out on flat thin specimens of a ferritic/pearlitic low strength steel and a acicular ferrite and bainite/martensite high strength steel, polarised at -1.05 and -2.00 V vs SCE, evidenced the effect of strain rate on the possibility of hydrogen to saturate the traps (Cabrini et al., 1998). At high strain rate (10–5 s–1 and higher) the hydrogen has not enough time to fulfill the traps as they are generated by the plastic straining. It was suggested that the equilibrium between the new traps generation by the external loading and their hydrogen saturation is achieved for strain rate lower than 10–6 s–1; at this strain rate the hydrogen permeation seems to be increased by the mechanical external stress (Cabrini et al., 1998, Bolzoni et al., 2001). At the lowest strain rate considered during tests (1.7∙10-7 s-1), critical potentials approach the value of thermodynamic hydrogen evolution. Thus, although at this potential the hydrogen adsorbed on surface is very low, a relevant quantity of hydrogen can penetrate in the metal lattice, favoured by the continuous straining. A slow and continuous straining is not only a necessary factor for crack initiation but it also play a fundamental role during propagation. If the specimen is maintained at constant load condition after the crack initiation (ISSR test results) the cracks stop to propagate. This behaviour was even observed at severe cathodic overprotection at -1.5 V vs SCE. Therefore, also the sub-critical propagation of the HE cracks requires the action of a continuous and slow straining. Nevertheless the presence of hydrogen in the steel increases the crack growth rate and decreases the fracture energy, as showed by means of J-integral tests carried out in absence or in presence of cathodic polarisation (Bolzoni et al., 2000). Hydrogen Embrittlement does not only depend on strain rate, but crack initiation takes place after a minimum critical strain is reached. In the monoaxial tensile SSR tests, HE phenomena take place only when plastic deformations were higher than those at necking beginning during tensile tests in air. Furthermore, the initiation of the cracks remains confined to the zone of the necking cone, where the highest plastic deformation is reached. Effect of microstructure on HE embrittlement susceptibility As the strain rate decreases, the critical value of potential increases in the range of -1.05 ÷ -0.85 V vs SCE depending on the steel microstructure and mechanical properties. Figures 17 and 18 show the critical potential as a function of strain rate and of the ultimate tensile strength of the steel. Rolled steels generally showed critical potential higher than the martensitic steels. All the examined steels produced by hot rolling and 6
controlled rolling show a similar behaviour at the lowest considered strain rate (1.7∙10-7 s-1), with critical potentials approaching the value of thermodynamic hydrogen evolution. Increasing the strain rate, the HE susceptibility of the steels seems to increase by the ultimate tensile strength of the steels. The critical potential approaches a minimum for the steel with UTS of about 700 MPa. The mechanical properties of these steels have been modulated by means of slight modifications in the chemical composition and in the parameters of thermo-mechanical process, mainly. The variations of microstructure and mechanical properties are described in the paragraph on materials in the Experimental section of this paper. The reduction of critical potential at low values of UTS is related with the pearlite-ferrite microstructures showed by rolling and controlled rolling steels, whereas the behaviour at higher strength is related to the recent grades of X80 and X100 steels with acicular microstructure. Koh et al. (2004) reports that ferrite-acicular microstructure has higher resistance to stress corrosion cracking in sour environments than ferrite-pearlite or ferrite-bainite microstructures. Many authors pointed out the better HIC and SSC resistance of acicular ferrite than polygonal ferrite and pearlite (Zhao et al., 2008, Park et al., 2008, Huang et al., 2011). A similar behavior can also be assumed under cathodic protection. In the ferritepearlite steels the HE cracks initiate in the pearlite bands and they do not propagate in the ferrite (Figure 19). The presence of hard constituents of un-tempered martensite decreases the cracking resistance (Liao and Lee, 1994, Park et al., 2008). The cracks are preferentially initiated in the martensite or bainite bands, but the crack growth is countered by the presence of acicular ferrite. The cracks appear very enlarged in the direction of loading (Figure 20), i.e. the rolling direction, with the HE propagation limited to the thickness of the pearlite bands, producing an almost ductile fracture surface. The HE resistance is strongly influenced by the amount of acicular ferrite respect to the bands of hard constituents like pearlite, bainite and martensite. As far as the HE resistance of rolled steels is concerned, it could be concluded that if tensile strength is produced by a very fine microstructure, HE resistance is not necessarily worsened. Thus, fine microstructure components as acicular ferrite, can counterweight the deleterious increase of tensile strength. However in steels with acicular ferrite HE susceptibility arises with the tensile strength owing to the effect of increasing contents of hard constituents, like un-tempered martensite. With respect to rolled steels, the tempered martensitic steels show critical potentials very scattered with a number of steels showing better behaviour than rolled steels (Figures 17-18). Literature data showed that quenching and tempering can improve the SSC resistance of HSLA steels, but exceptions were observed (Staehle, 1969). Carneiro et al. (2003) reported that a refined homogeneous quenched and tempered bainite/martensite microstructure simultaneously improved HIC and SSCC resistance. These steels showed higher SSCC resistance than both rolled and normalized steels with ferritic/pearlitic microstructures. Albarran et al. (1999) confirmed that quenching and tempering provides an optimum SSCC resistance of API X-80 steels. The improved SSCC resistance of the quenched and tempered microstructure can be attributed to its homogeneous bainite/martensite microstructure, which made difficult crack propagation (Wang and Staehle 1972, Kobayashi et al., 1982). Because banded structure provides high hydrogen trap density can enhance crack propagation with respect to homogenous microstructures (Carneiro et al., 2003, Albarran et al., 2003). Thus homogeneous microstructures with dispersed and very fine traps can distribute hydrogen on high number of sites, so that it is less probable to reach the critical concentration. Figure 21 compares the behaviour of TM3, TM4 and TM5 steels which have same tensile strength and close chemical compositions with only different nickel contents. They show different resistance to HE. The TM3 steel is less resistant than the other two steels which perform better than steels with banded ferrite microstructures. The resistance to HE may be related to the morphology of precipitates, mainly carbides. A recent paper emphasizes the role of martensite-austenite phase (Park, 2008). The Figure 22 shows the microstructures of the steels after picral etching in order to evidence carbide distribution. TM3 steel has a microstructure characterized by evident lath martensite whereas TM4 steel and TM5 steel have fully tempered martensite. The steels show different distribution of carbides after picral etching (Figure 21). TM4 steel shows very fine precipitate distribution. On the other hand, TM3 steel shows few particles of coarse precipitates. 7
HE cracking occurs on grain boundaries of precipitates along paths where they are concentrated (Figure 23). At this location, the interfacial decohesion can be favored by the combination of a hydrogen build-up coupled with significant internal stresses as outlined by Albarran et al. (1999). The precipitate interface can act as hydrogen trap. A fine dispersion of carbides raises the extension of interface so that low local hydrogen contents are only achieved (Tau et al., 1996), increasing HE resistance of the steel. CONCLUSIONS The behaviour of different HSLA linepipe steels towards the Hydrogen Embrittlement under cathodic protection were evaluated. The results confirm the immunity of all steels under static loading conditions. HE phenomena were never noticed up to yield strength, even at very negative cathodic protection potentials. They were only observed under slow and continuous plastic straining typical of the SSR tests. The presence of a slow and continuous straining is a necessary factor both for crack initiation and propagation. HE phenomena occurs at potential less noble than a critical potential, which depends on strain rate, steel microstructure and its mechanical properties. Quenched and tempered steels with high tempered martensite and very fine precipitation distribution show a better behavior than the rolled steels with banded microstructure, exhibiting the highest HE resistance. The resistance of steels characterized by banded microstructures increases with the ultimate tensile strength until about 700 MPa, than decreases. The increase of ultimate tensile strength by changing the steel microstructure, from hot rolling with coarse ferrite/pearlite to fine ferrite/pearlite microstructure, and finally very fine acicular ferrite arises the resistance of steel to HE. Further increase of ultimate tensile strength related with un-tempered martensite inside the acicular ferrite based microstructure produce a worsening of resistance. References Albarran J.L., Martinez L., Lopez H.F., Corrosion Sci. 41 (1999) 1037-1049. Albarran J.L. et al., Stress corrosion cracking response of microalloyed pipeline steels exposed to pressurized sour gas environments. Corrosion 2003, paper 3447, March 16–20, San Diego, CA. ASTM D1141-90 Standard Practice for the Preparation of Substitute Ocean Water. ASTM E399-90 Standard Test Method for Linear-Elastic Plane-Strain Fracture Toughness KIc of Metallic Materials. Barsanti L., Bolzoni F.M., Cabrini M., Pastore T., Spinelli C., “Hydrogen Embrittlement resistance of X100 steels for long distance high pressure pipelines”, Second International Conference on Environmental-Induced Cracking of Metals EICM04, September 19-23, 2004, The Banff Centre, Banff, Canada. Beidoknti B., A. Dolati, A.H. Koukabi, Mater. Sci. Eng. A507 (2009) 167–173. Bolzoni F.M., M.Cabrini, P.Pedeferri, C.Spinelli, “Evaluation of Hydrogen Embrittlement Resistance of Microalloyed Steels by means of J-integral Curve”, Proceedings of EUROCORR2000, Institute of Corrosion, London, 2000. 8
Bolzoni F., M.Cabrini, C. Spinelli, “Hydrogen diffusion and hydrogen embrittlement behaviour of two high strength pipeline steels”, Eurocorr 2001, Riva del Garda, Proceedings on CD-ROM, AIM, Milano, pag. 1-10. ISBN 88-85298-41-9. Cabrini M., S.Maffi, G.Razzini, “Evaluation of Hydrogen Embrittlement behaviour by means permeation current measure in slow strain rate conditions of a micro-alloyed steel”, Materials Science Forum, UK (ed. Trans Tech Publications) p.1245 (1998). Cabrini M., F.Bolzoni, G.Razzini, E.Sinigaglia, 14th ICC, Paper 30 (1999). Cabrini M., L.Migliardi, T.Pastore, C.Spinelli, ”Effect Of Cathodic Potential And Strain Rate on Hydrogen Embrittlement of HSLA Steels”, in Hydrogen Effectes on Materials Behaviour and Corrosion Deformation Interaction, Edited by N. R. Moody, A.W:Thompson, R.E.Ricker, G.W.Was and R.H.Jones, TMS (The Mineral, Metals & Materials Society) (2003) pp 979-988. Cabrini M., O.Cogliati, S.Maffi, “Effetto della microstruttura sulla diffusione dell’idrogeno in acciai al carbonio per pipeline”, La metallurgia italiana, N. 3 (2003) pp.13-20. Cabrini M., G.D’Urso, T.Pastore, “Evaluation of the resistance to hydrogen embrittlement by slow bending test”, Second International Conference on Environmental-Induced Cracking of Metals EICM04, September 1923, 2004, The Banff Centre, Banff, Canada. Cabrini M., T.Pastore, L.Barsanti, C.Spinelli, “Effect of microstructure on the behaviour of HSLA steels to hydrogen embrittlement under cathodic protection”, Second International Conference on EnvironmentalInduced Cracking of Metals EICM04, September 19-23, 2004, The Banff Centre, Banff, Canada. Carneiro R.A. et al. / Materials Science and Engineering A357 (2003) 104-110. Carter C.S. and M.V. Hyatt, "Review of Stress Corrosion Cracking in Low Alloy Steels with Yield Strength Below 150 ksi", 1973, in: R.W. Staehle, J.Hochman, R.D. McCrigth, J.E. Slater (Eds.), Stress Corrosion Cracking and Hydrogen Embrittlement of Iron Base Alloys, NACE, Houston TX, 1977, pp. 524-600. Chiovelli S.C., Dorling D.V., Glover A.G., Horsley D.J., Oil Gas J 1994;92(11):91–100. Cigada A., Pastore T., Re G., Rondelli G., Vicentini B., Stress corrosion cracking in seawater of HSLA steel type API 5L X65, Metals/materials technology series, ASM 8520-012, 1985. Cigada A., T.Pastore, G.Franzoso, A. Kopliku, M.Carenza, “SCC of high strength steels in deep seawater under cathodic protection”, Offshore Tecnology Conference, paper 5156, 1992, ISOPE-92. Delanty B., O’Beime J., Oil&Gas Journal, June 15, 1992, p.39. Demofonti G., M.Cabrini, F.Marchesani, C.M.Spinelli, “Eni Tap Project Mechanical Damage and Environmental Assisted Cracking -Full Scale Methodology Overview”, Proceedings of New Developments on Metallurgy and Applications of High Strength Steels Conference, May 26 - 28, 2008, Buenos Aires, Argentina. Det Norske Veritas RP B401, Cathodic Protection Design, 1986. Engell H.J., Forchhammer P., Uber die veranderung des pH-wertes an metalloberflachen bei kathodischem 9
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Punter A., A.T. Fikker, G. Venstoen, Material Performance 31 p.24 (1992). Razzini G., Cabrini M., Maffi S., Peraldo Bicelli L., Musatti G., “Effect of the heat-affected zones on hydrogen permeation and embrittlement of low-carbon steels”, Materials Science Forum Vols. 286–292 (1998) pp 1257– 1266, Trans Tech Publications, Switzerland. Razzini G., M.Cabrini, S.Maffi, G.Mussati and L.Peraldo Bicelli, Corrosion Science 41 p.203 (1999). Rodrigues P.C.M. et al. /Materials Science and Engineering A283 (2000) 136–143. Rondelli G., Re G., Cigada A., Vicentini B., Pastore T., Stress corrosion cracking in sea water of a steel for offshore applications, The American society of Mechanical Engineers, 1987. Saenz de Santa Maria M., R.P.M. Procter, Environmental cracking (corrosion fatigue and hydrogen embrittlement) X-70 linepipe steel. In: Fatigue and crack growth in offshore structures, C137/86. London: IMechE; (1986) p. 101–8. Seo D.H., S.H.Cho, C.M.Kim,J.Y.Joo, K.B.Kang, Microstructure and mechanical properties of X80/API 5L X100 plates and pipes. In proc. The 17th International Offshore and Polar Engineering Conference, Lisbone, Portugal, 2007. Shanmugametal S., Materials Science and Engineering A 460–461 (2007) 335–343. Shipilov S.A., Critical assessment of the rule of cathodic protection in pipeline integrity and reliability. In: Flewitt PEJ et al., editors. Engineering structural integrity assessment: need and provision. Sheffield: EMAS; 2002. P. 155-62. Shipilov SA. Corrosion fatigue. In: Varvani-Farahani A, editor. Advances in fatigue, fracture and damage assessment of materials. Southampton: WIT Press; 2005. Shipilov S.A., I.Le May, Engineering Failure Analysis 13 (2006) 1159-1176. Staehle A., J.Forty, D.Van Rooyen (Eds.), Fundamentals Aspects of Stress Corrosion Cracking. NACE, Houston, TX, 1969, p. 361. Tanguy B., T.T.Luu, G.Perrin, A.Pineau, J.Besson, Plastic and damage behaviour of high strength API5L X100 pipeline steel: experiments and modelling. International Journal of pressure Vessel and Piping (2005); 85:32235. Tau L., S.L.I. Chan, C.S.Shin, Corrosion Science, 38 (1996) 2049-2060. Trasatti S.P.et al., Materials and Corrosion, 2005, 56 (2), 111-17. Wang M.T., R.W. Staehle, Proceedings of the Hydrogen in Metals: International Conference, Paris, May /June 1972, pp. 342. Wei R.P., Novak S.R., “Interlaboratory evaluation of KISCC and da/dt determination procedures for high strength steels”, Journal of testing and evaluation, 15 (1) 38, 1987. Zhao M.-C. et al., Materials Science and Engineering A 478 (2008) 43–47 11
Appendix - Evaluation of alkalinization on metallic surface under cathodic protection The concentration of OH- on steel surface has been calculated by mass balance flux at metal surface in stationary conditions (Engell and Forchhammer, 1965), by considering the diffusion flux and Faraday’s Law. By assuming that salinity of solution is given by chemical ions not directly involved in the reactions, the electrophoretic contribution of mass transport flux can be ignored. At the hydrogen equilibrium potential the main cathodic reaction in aerated solution under cathodic polarization is 𝑂$ + 4𝑒 ( + 2𝐻$ 𝑂 → 4𝑂𝐻 ( (7) Therefore OH ions production rate is equal to oxygen limiting current density: 𝑖FGH,IJ = 4 ∙ F ∙ 𝑘IJ ∙ 𝐶N,IJ (8) where F is the Faraday constant, kO2 is the mass transport coefficient of oxygen in water and CB,O2 the bulk oxygen concentration. The diffusion flux of ionic species from the surface can be expressed by relation 𝑖O,G = F ∙ 𝑘G ∙ P𝐶Q,G − 𝐶2,G R (10) where ki is the mass transport coefficient of the ion in water, C0 and CB are respectively the concentrations on surface and in bulk solution. OH- ions produced by cathodic reaction are consumed by water formation reaction, negligible if H + concentration is low, such as in neutral-alkaline solutions, and by calcium carbonate precipitation above saturation pH, according to: 𝐶𝑎TT + 𝑂𝐻 ( + 𝐻𝐶𝑂U( → 𝐶𝑎𝐶𝑂U (11) Case 1: non-scaling water In neutral solution without precipitation of carbonates, it can be assumed that the production of OH-, equal to the oxygen limiting current, is balanced by the diffusion from the electrode: 4𝑘IJ ∙ 𝐶N,IJ = 𝑘VW7 ∙ P𝐶2,VW7 − 𝐶Q,VW7 R (14) The mass transport coefficient can be calculated as a function of Reynolds (Re), Schmidt (Sc) and Sherwood (Sh) adimensional numbers by using literature equations (Hara et al., 2000), generally expressed as b Sh = cost ∙ 𝑅𝑒 _ ∙ 𝑆𝑐 cU (15) With Sherwood number and Schmidt number respectively defined by Shd =
ef ∙g hf
;
i
Scd = h
f
(16)
where ν is the kinematic viscosity and l the characteristic length. Re is a function of fluidodynamic conditions only, whereas Sh and Sc depend on chemical species because they depend on diffusion coefficient. By substituting (15-16) in (14) $c
$c
4𝐷IJU ∙ 𝐶N,IJ = 𝐷VWU7 ∙ P𝐶2,VW7 − 𝐶Q,VW7 R thus lhmJ
𝐶2,VW7 = 𝐶Q,VW7 + kh
no7
p
$c U
∙ 𝐶N,IJ
(18)
(19)
Case 2: scaling water If calcium carbonate can separate at the surface, the pH on the electrode is reduced by buffer capacity of the solution. The alkalinity modifies the bicarbonate/carbonate equilibrium and produces precipitation of calcium carbonate 𝐻$ 𝑂 ⇄ 𝐻 T + 𝑂𝐻 ( (20) ( (( T 𝐻𝐶𝑂U ⇄ 𝐶𝑂U + 𝐻 (21) (( TT 𝐶𝑎𝐶𝑂U ⇄ 𝐶𝑂U + 𝐶𝑎 (22) 12
The equilibrium of reaction (11) can be expressed by rJ b = 7 ∙u r ∙r u ∙u 7 s
t
;vmw
vxyy
m;
(23)
where Kw, K2 and Ks are respectively the water ionic product, the second dissociation constant of carbonic acid and the solubility product of calcium carbonate. The flux of OH- ions produced at cathode is balanced by the flux of ions diffusing from surface to bulk solution and the consumption of ions (iR,i) due to reaction (11) 𝑖FGH,IJ = 𝑖O,VW7 + 𝑖z,VW7 (24) By considering that the consumption of hydroxide ions is equal to that of calcium and bicarbonate ions 𝑖z,VW7 = 𝑖z,{|yy = 𝑖z,0uIw7 (25) For the mass balance of calcium and bicarbonate ions 𝑖O,{|yy = 𝑖z,VW7 (26) 𝑖O,0uIw7 = 𝑖z,VW7 (27) Finally, the hydroxide concentration can be calculated from equation (10) by the following relations and (23) $c
$c
$c
$c 𝐷VWU7
$c U 𝐷0uI 7 w
4𝐷IJU ∙ 𝐶N,IJ − 𝐷VWU7 ∙ P𝐶2,VW7 − 𝐶Q,VW7 R = 𝐷{|Uyy ∙ P𝐶Q,{|yy − 𝐶2,{|yy R
(30)
(31)
However the precipitation can only occur from supersatured solutions above saturation pH: r ∙r 𝐶2,0uIw7 ∙ 𝐶2,u*yy ∙ 𝐶2,I07 = 𝛼 ∙ sr t (𝛼 ≥ 1)
(33)
$c 4𝐷IJU
∙ 𝐶N,IJ −
∙ P𝐶2,VW7 − 𝐶Q,VW7 R =
J
∙ P𝐶Q,0uIw7 − 𝐶2,0uIw7 R
Figure 24 shows the variation of pH on cathode surface as a function of oxygen concentration estimated, without any carbonate precipitation and with carbonate precipitation calculated by considering a value of α equal to 3.
13
Table 1: Chemical composition, mechanical properties and microstructure of the tested HSLA steels. Steel
C
Mn Si
P
S
Cr
Ni Mo Nb
Cu
TYS UTS Micro- (MPa) (MPa) structure
Hot rolled steel (longitudinal welded pipelines) FP1 0.22 1.35 0.24 0.012 0.024
