FAILURE OF LOW ALLOY STEELS
Since the majority of the bolts used in undersea connections are made of low alloy steels with carbon concentrations in the range of 0.3-0.4 wt.%, this appendix will focus on failure modes for this type of material. Hydrogen embrittlement in Ni-based alloys is briefly discussed since a limited number of bolts are made from nickel based superalloys.
The general characteristics of low alloy steels that relate to fracture response are reviewed to provide background information for a discussion of hydrogen embrittlement. These materials undergo a transition from brittle to ductile failure with increasing temperature. The temperature at which this occurs, referred to as the ductile-to-brittle transition temperature (DBTT), is usually defined as the midpoint in a range of temperatures, spanning approximately 50°, during which the fracture mode transitions from 100 percent brittle to 100 percent ductile. The transition temperature is most commonly determined through Charpy testing, although a notched fracture test in which the sample is loaded in tension can be used. The latter type of test reports a fracture toughness value for crack propagation (KIC), whereas the Charpy test simply gives integrated fracture energy. The transition temperature is usually well below room temperature for steels that have been tempered in the range of 600°C to 700°C and then rapidly cooled. For applications such as the one considered in this appendix and for steels heat treated according to the current standards, the fasteners should fail in a ductile manner.
Ductile failure in low alloy steels is generally associated with a high fracture energy, which implies that significant stress must be applied to the material for it to occur. The mode of fracture is often described as micro-void nucleation and coalescence and evolves in the following way. As the sample is loaded in tension, the interface between the steel matrix and large inclusions, the most common of which are manganese sulfides, separates and creates microvoids that grow with increasing applied stress. This growth reduces the cross-sectional area over which the stress is applied. When the stress reaches a critical value the material between these large voids rips apart and failure occurs. The fracture path between these large voids is often populated with much smaller voids that form around the carbides produced during tempering and is usually transgranular.
Low energy ductile fracture can be observed in some situations as a result of material processing. This low energy fracture can occur as a result of a high density of sulfides and alignment of the sulfides in a particular direction in the material that is also parallel to the fracture path. It can also result from a process referred to as overheating which causes precipitation of the sulfides on the grain boundaries.1,2 In the latter case the ductile fracture mode is intergranular.
Brittle fracture occurs with much lower energy absorption. The general model for this type of fracture, often referred to as the Griffith model, is that a microcrack is nucleated at a second phase particle or defect and then, once it reaches a critical size, propagates rapidly across the sample.3 In steels that have been quenched from the tempering treatment, the most common sites of nucleation are cracked carbides or cracked inclusions.4 However, if the sample is not rapidly quenched and the steel contains impurity elements such as phosphorus or tin at levels above approximately 100 wt ppm, the grain boundaries can become weak and serve as the sites of crack nucleation and propagation.5 In this case, brittle fracture can be observed at much higher temperatures than expected. This phenomenon is referred to as temper embrittlement.
Unpredicted brittle fracture can also be caused by a specific environment. Many types of environments can be listed that can cause low energy fracture in steel, but here the discussion is limited to hydrogen embrittlement. Hydrogen embrittlement
1 T.B. Cox and J.R. Low, An investigation of the plastic fracture of AISI 4340 and 18 Nickel-200 grade maraging steels, Metallurgical Transactions 5(6):1457-1470, 1974.
2 A.M. Ritter and C.L. Briant, “The Effect of Second-Phase Particles on Fracture in Engineering Alloys,” pp. 59-123 in Embrittlement of Engineering Alloys (C.L. Briant and S.K. Banerji, eds.), Academic Press, New York, N.Y., 1983.
3 A.A. Griffith, Philosophical Transactions of the Royal Society A 221:163, 1920.
4 C.J. McMahon, Jr., and M. Cohen, Initiation of cleavage in polycrystalline iron, Acta Metallurgica 13:591, 1965.
5 C. L. Briant and S.K. Banerji, Intergranular failure in steel: The role of grain-boundary composition, International Metals Reviews 23:164, 1978.
has, in general, been described by three broad mechanisms which are important because they point to certain dependencies, triggers, and drivers. Hydrogen induced decohesion (HEDE) relates to a lowering of the metal-metal bond strength and Griffiths fracture toughness in response to hydrogen. HEDE can be exacerbated by the impurity segregation described above. Hydrogen induced local plasticity relates to hydrogen enhanced deformation that results in a variety of effects such as enhanced plastic slip, greater dislocation planarity and slip band cracking, and large local accumulations of dislocation bands such that detrimental pile-ups occur at interfaces like grain boundaries which in turn trigger intergranular fracture. The third mechanism involves hydriding where a metal-hydrogen phase forms that has a low intrinsic fracture toughness. The first two mechanisms are the most likely operative in ferrous and nickel base alloys used as oil and gas marine connectors. It should be noted that nickel and nickel alloys such as Ni-Cu may hydride albeit not at typical cathodic potentials used in marine service.
The source of hydrogen in marine applications can arise from a combination of factors. These include processing such as pickling, coatings, or post-coating bake-outs at the OEM, long term field exposure where water is reduced electrochemically by galvanic coupling, impressed current or sacrificial anode based cathodic polarization, or anoxic freely corroding conditions at high sensible pressures.
- The crack growth is time dependent. A crack can propagate in a stable fashion, once it reaches a critical threshold stress intensity, and continue until it reaches a length such that, for a given applied stress, rapid fracture takes over. Thus, it is not a failure that will necessarily occur immediately after a part is loaded or put into service. It is generally assumed that this
6 R.P. Gangloff and R.P. Wei, Gaseous hydrogen embrittlement of high strength steels, Metallurgical Transactions A 8:1043-1053, 1977.
7 J.R. Pickens, J.R. Gordon, and J.A.S. Green, The effect of loading mode on the stress-corrosion cracking of aluminum alloy 5083, Metallurgical Transactions A 14:925, 1983.
8 H.K.D.H. Bhadeshia, “Extremely Strong Steels—The Mechanism and Prevention of Hydrogen Embrittlement,” Lecture, AISTech 2017, Association for Iron and Steel, May 8, 2017.
9 Z.D. Harris, J.D. Dolph, G.L. Pioszak, B.C.R. Troconis, J.R. Scully, J.T. Burns, The effect of microstructural variation on the hydrogen environment-assisted cracking of Monel K-500, Metallurgical and Materials Transactions A 47:3488, 2016.
10 R.P. Gangloff, H. Ha, J. Burns, and J.R. Scully, Measurement and modeling of hydrogen environment-assisted cracking in Monel K-500, Metallurgical and Materials Transactions A 45(9):3814-3834, 2014.
11 J.H. Ai, H.M. Ha, R.P. Gangloff, and J.R. Scully, Hydrogen diffusion and trapping in a precipitation-hardened nickel-copper-aluminum alloy Monel K-500 (UNS N05500), Acta Materialia 61(9):3186-3199.
- time dependence occurs because hydrogen must continue to diffuse to the crack tip as the crack propagates. Hence the hydrogen diffusion coefficient at the crack tip is critical. This parameter is subject to modifications by the details of the plastic and fracture process zones; hydrogen trapping and concentration dependent diffusion are critical factors.
- The threshold stress intensity for cracking is reduced and the crack growth rate is increased by the diffusible hydrogen concentration to a different degree for a given material. However, the sensitivity of a material to a given hydrogen content depends on many material factors such as strength, Griffiths fracture toughness (in a decohesion model), and the potency factor for hydrogen which may in turn be related to grain boundary structural or chemical factors such as interface impurity and metalloid content.
- A critical hydrogen content for fracture is often observed. A subtle detail is that hydrogen may affect ductile cracking processes, although this type of fracture is a much higher energy fracture than brittle failures caused by hydrogen embrittlement. At greater hydrogen contents these fracture modes may transition with appropriate tensile stress, microstructure and local hydrogen content to increasing brittle modes such as intergranular and lath boundary cracking. Hence the notion of a threshold may correspond to any detectable embrittlement or mode transition.
- Cracking and susceptibility may be strain rate dependent and affected by dynamic plastic strain during or after hydrogen charging. The mechanistic effects of strain are debated and may alter surface effects such as hydrogen production and uptake (e.g. enhanced by film rupture) versus internal effects such time for diffusion controlled transport to the fracture process zone, dynamic trap creation, and dislocation transport from the tip into the fracture process zone.
- The rate of crack growth is temperature dependent. For steels, a maximum in crack growth rate is often near or slightly below room temperature. This observation is often related to a balance between trapping and diffusion where detrapping and outgassing occur at high temperatures and transport is very slow at cold temperatures.
- A tensile stress is required to cause the failure. Hydrogen diffuses to regions of high tensile stress in a material, such as those at a crack loaded in tension, and concentrates there.
- The fracture mode can be either transgranular or intergranular. The details of the fracture mode depend on the concentration of hydrogen and also the susceptibility of the grain boundaries to fracture as discussed below.
- Traps in the material, such as precipitates or dislocations, can increase the amount of hydrogen required to cause failure. Some strong traps effectively pin the hydrogen and keep it from diffusing to areas of high tensile stress
- in a closed system. This only works for low crack tip stress fields. This can only delay cracking in an open system as trap filling and saturation will allow hydrogen to eventually accumulate. Reversible traps of intermediate and low strength can impede diffusion even in an open system where a slower effective diffusivity can be affected by intermediate strength cracks to limit stage II steady state crack growth rates even in open systems.12
In the next section, detailed aspects of HAC in steels that relate to the topic of this appendix are discussed. In particular, these examples will show that the hydrogen concentration in the sample, and particularly at the crack tip, plays a critical role in determining the degree of embrittlement. If this concentration is sufficiently high, most steels will be susceptible to hydrogen embrittlement.
General HE Susceptibility
Of all the steel types used in subsea connectors, many are susceptible to hydrogen embrittlement in seawater at free corrosion potentials, with a common denominator being strength level in excess of about 80-110 ksi (560 to 800 MPa). These include bainitic, martensitic and martensitic and aged steels with the precise details of embrittlement depending on microstructure, interface cleanliness and strength. This behavior is shown in Figure K.1 where exposure of pre-cracked specimens to 3.5 percent NaCl is shown to lower the fracture toughness of low alloy steels, that range in strength from 115 to over 260 ksi (800 to over 1800 MPa), from over 55 ksi(in)1/2 to 9-18 ksi(in)1/2 (60 MPa(m)1/2 to 10-20 MPa(m)1/2).13 It should be noted that similar behavior14 occurs in quenched and tempered reactor steels.15
The root cause of SCC of steels in seawater has been shown to be diffusible hydrogen concentration within the alloy regardless of source.16 The stress intensity dependence of subcritical crack growth rate produced in a tempered martensitic
12 Open systems have continual hydrogen production while closed systems does not.
13 R.P. Gangloff, A Review and Analysis of the Threshold for HEE of Steel, in Corrosion Prevention and Control, Proceedings of the 33rd Sagamore Army Materials Research Conference (S. Isserow, ed.), U.S. Army Materials Technology Laboratory, Watertown, Mass., 1986.
14 An exception to this well-established trend is in the case of extreme sour gas charging seen in the NACE MR 0175 test solution. In so-called, “sulfide induced stress corrosion cracking,” cracking may be controlled by inclusion density and shape and be independent of yield strength. Only in this highly specialized case is the strong yield strength correlation is absent.
15 W.E. Erwin and J.G. Kerr, “The Use of Quenched and Tempered 2 1/4Cr-1Mo Steel for Thick Wall Reactor Vessels in Petroleum Refinery Processes: An Interpretive Review of 25 Years of Research and Application,” Welding Research Council, 1982.
16 R.P. Gangloff, “Hydrogen Assisted Cracking of High Strength Alloys,” pp. 31-101 in Comprehensive Structural Integrity (I. Milne, R.O. Ritchie, B.L. Karihaloo, eds.), Elsevier/Pergamon, Amsterdam, Boston, 2003.
steel exposed in three separate environments that produce atomic hydrogen at the crack tip are shown in Figure K.2.17 Here H2S and H2 gas are more severe than 3.5 percent NaCl and an overall toughness decreases to the range of 10-20 MPa (m)1/2 is observed at applied stress intensities above these thresholds, a stage I region is typically followed by a stage II region of crack growth rate that strongly depends on the environment and potential and, in turn, the hydrogen concentration at the crack tip. Thus the H2S environment is more severe than seawater because the hydrogen
17 G.E. Kerns, M.T. Wang, and R.W. Staehle, pp. 700-733 in Stress Corrosion Cracking and Hydrogen Embrittlement of Iron Base Alloys (R.W. Staehle, J. Hochman, D.R. McCright, and J.E. Slater, eds.), National Association of Corrosion Engineers, Houston, Tex., 1973.
concentration is likely greater. This extreme susceptibility is not encountered at lower strength levels as shown in Figure K.2 but immunity is not observed.
It is notable that while at very high strength levels, maraging steels are also susceptible to stress corrosion cracking in seawater, they are less susceptible than quenched and tempered 4340 steel at the same strength level due to differences in the microstructure.18,19 Vacuum melted, quenched, and tempered steels with low amounts of trace elements compare even more favorably with maraging steels in resisting stress corrosion at high-strength levels. For example, modern clean UHSS such as Aermet 100TM do not crack from prior austenite grain boundaries but instead exhibit cracking at lath interfaces.20 Unfortunately the drop in toughness with hydrogen content even with a carefully controlled microstructure is still significant.
Dependence on Applied Potential
Electrode potential is of great importance in the behavior of steels because these alloys are often coated with a sacrificial anodic material such as zinc for corrosion protection which creates a galvanic couple and/or are cathodically protected by either sacrificial anodes or impressed current cathodic protection systems. One of the key factors in the type and severity of hydrogen induced cracking of high strength steels is the electrode potential. This is shown in Figures 2.10 through 2.12.21,22 Brown first reported the effect of applied potential on crack growth rate in high strength steels.23 Crack growth rate was found to increase at both highly anodic and highly cathodic applied potentials with a reduced susceptibility at intermediate
18 D.G. Enos and J.R. Scully, A critical-strain criterion for hydrogen embrittlement of cold-drawn, ultrafine pearlitic steel, Metallurgical and Materials Transactions A 33:1151-1166, 2002.
19 R.N. Parkins, A.J. Markworth, J.H. Holbrook, and R.R. Fessler, Hydrogen gas evolution from cathodically protected surfaces, Paper 47 in Corrosion 84, National Association of Corrosion Engineers, Houston, Tex., 1984.
20 Y. Lee and R.P. Gangloff, Measurement and modeling of hydrogen environment-assisted cracking of ultra-high-strength steel, Metallurgical and Materials Transactions A 38:2174-2190, 2007.
21 B.A. Kehler and J.R. Scully, Predicting the effect of applied potential on crack tip hydrogen concentration in low-alloy martensitic steels, Corrosion 64:465-477, 2008.
22 B.F. Brown, Stress-Corrosion Cracking in High Strength Steels and in Titanium and Aluminum Alloys, Naval Research Laboratory, Washington, D.C., 1972, p. 3.
23 C.T. Fujii, “Stress-Corrosion Cracking Properties of 17-4 PH Steel,” p. 430 in Stress Corrosion—New Approaches: A Symposium Presented at the Seventy-Eighth Annual Meeting of the American Society for Testing and Materials (H.L. Craig, ed.), American Society for Testing and Materials, Philadelphia, Pa., 1976.
potentials.24 This behavior was seen in K92580 and 300M and ESR 4340 steels.25,26 PH 15-5 stainless steel,27 PH 13-8,28 maraging steels,29 eutectoid cold drawn and heat treated steels,30 C1045 steel,31,32 PH 17-4, 9-4-45 and 4340 high strength steels and AISI 1080 steel,33 and in modern variants of UHSS such as Aermet 100. It is clear that a broad range of different high strength steel alloys exhibit this behavior.
This effect is shown in the potential dependent fracture toughness data in Figure K.3. HEAC susceptibility is strongly potential dependent on both cathodic and anodic potentials in many UHSS alloys; applied potential (EApp) affects both threshold stress intensity (KTH) and stage-II subcritical crack growth rate (da/dtII). Figure 2.11 shows da/dtII in AerMet®100 as a function of EApp in 0.6 M NaCl and the open circuit potential for various coatings in chloride solutions. The root cause is the diffusible hydrogen concentration at the crack tip developed as a function of external applied potential and crack chemistry as examined and modeled by Turnbull and modeled/predicted by Kehler.
Figure K.4 also indicates the predicted diffusible hydrogen concentration at the crack tip as a function of applied potential. The cause of HE is the establishment of a high CH,diff at the crack tip governed by chemical and electrochemical factors, as well as crack tip metallurgy, where CH,diff is the diffusible hydrogen content or
24 B.F. Brown, “Stress Corrosion Cracking of High Strength Steels,” p. 471 in The Theory of Stress Corrosion Cracking in Alloys (J.C. Scully, ed.), North Atlantic Treaty Organization, Scientific Affairs Division, Brussels, Belgium, 1971.
25 E.U. Lee, H. Sanders, and B. Sarkar, “Stress Corrosion Cracking of High Strength Steels,” Proceedings of the Tri-Service Conference on Corrosion (J.V. Kelley and B. Placzankis, eds), U.S. Army Research Laboratory, Aberdeen, Md., 2000.
26 P.F. Buckley, R.H. Brown, B. Placzankis, and J. Beatty, Paper 547 in Corrosion 94, National Association of Corrosion Engineers, Houston, Tex., 2004.
27 K.B. Das, W.G. Smith, R.W. Finger, and J.N. Master, “Hydrogen Embrittlement of Cathodically Protected 15-5PH Stainless Steel,” in Proceedings of the Second International Congress of Hydrogen in Metals, International Association of Hydrogen Energy, Pergamon, Paris, France, 1977.
28 P.S. Tyler, M. Levy, and L. Raymond, Investigation of the conditions for crack propagation and arrest under cathodic polarization by rising step load bend testing, Corrosion 47:82-87, 1991.
29 D.P. Dautovich and S. Floreen, “The Stress Corrosion and Hydrogen Embrittlement Behavior of Maraging Steels,” pp. 1210, 1215 in Stress Corrosion Cracking and Hydrogen Embrittlement of Iron Base Alloys (R.W. Staehle, ed.), National Association of Corrosion Engineers, Houston, Tex., 1977.
30 V.S. Galvez, L. Caballero, and M. Elices, “The Effect of Strain Rate on the Stress Corrosion Cracking of Steels for Prestressing Concrete,” p. 603 in Laboratory Corrosion Tests and Standards: A Symposium by ASTM Committee G-1 on Corrosion of Metals (G.S. Haynes and R. Baboian, eds.), ASTM, Philadelphia, Pa,, 1985.
31 C.F. Barth, E.A. Steigerwald, and A.R. Trojano, Hydrogen permeability and delayed failure of polarized martenstic steels, Corrosion 25:353, 1969.
32 C.F. Barth and A.R. Troiano, Cathodic protection and hydrogen in stress corrosion cracking, Corrosion 28:259, 1972.
33 R.M. Schroeder and I.L. Muller, Stress corrosion cracking and hydrogen embrittlement susceptibility of an eutectoid steel employed in prestressed concrete, Corrosion Science 45:1969-1983, 2003.
concentration (moles/cm3 or wt. ppm) consisting of the sum of both the perfect lattice hydrogen for a given hydrogen fugacity obtained through the cathodic hydrogen overpotential and the hydrogen concentration associated with weak and intermediate reversible hydrogen traps. The hydrogen occluded at these traps is in equilibrium with hydrogen in lattice sites. The diffusible hydrogen is mobile and in general can repartition to the tensile tri-axial stress field of the crack tip.
Figure K.5 details the strong effect of hydrogen concentration on threshold stress intensity in H2 gas and in aqueous solutions. Unlike the IHE situation that involves a fixed amount of hydrogen in a closed system, a continuous supply of
hydrogen at the crack tip in such an open system presents a greater challenge to mitigation of TG HEE. Thus, degradation of fracture toughness is a strong function of applied electrode potential which enables the continual supply of hydrogen in marine environments.
Under cathodic polarization, the crack becomes increasingly alkaline relative to the bulk as a result of proton discharge and water reduction. In addition, the crack tip potential is shifted to more positive potentials due to ohmic voltage drop. Therefore, the overpotential for hydrogen production may actually be smaller at the crack tip than on the boldly exposed surfaces in the cathodic case. Thus, the hydrogen overpotential is lower and the pH is high and hydrogen uptake is lower. This is confirmed through visual inspection of the fracture surface. When the overpotential for hydrogen production is greater on the bulk surfaces than at the crack tip, the crack grows faster on the sample edges than at mid thickness. That is, the sample edge corresponds to higher values of CH,diff, while the sample center corresponds to lower CH,diff values. Therefore, cathodic polarization may lead to greater hydrogen uptake at less occluded sites, with faster near surface
cracking even though the stress intensity and hydrostatic stress may be lower at such locations. It should be further noted that anodic potential produce hydrogen embrittlement in high strength materials and not a switch in mechanism towards stress corrosion cracking by anodic dissolution.
It is readily shown that acidification and crack tip hydrogen uptake occurs in anodically polarized high strength steels. When normalizing to hydrogen content such as CH,diff, it can be seen that in cases of both anodic and cathodic polarization cracking is similar. Also, hydrogen absorption characteristics could be markedly different on an actively corroding or plastically strained surfaces compared to unstrained metal in corroded condition or covered with calcareous deposits.34,35
34 J.R. Scully, M.J. Cieslak, and J.A. Vandenavyle, Hydrogen embrittlement behavior of palladium modified PH 13-8 Mo stainless steel as a function of age hardening, Scripta Metallurgica et Materialia 31:125-130, 1994.
35 J.R. Scully and P.J. Moran, The influence of strain on hydrogen entry and transport in a high strength steel in sodium chloride solution, Journal of the Electrochemical Society 135:1337-1348, 1988.
The above analysis points out that the level of cathodic or anodic polarization, affected by impressed current cathodic protection level, galvanic coupling material or corrosion protection coating is extremely important. Consequently, the danger of hydrogen embrittlement must be kept in mind during the application of cathodic protection to alloys, such as martensitic stainless and other high strength steels, known to be susceptible to hydrogen embrittlement. For instance, cathodic protection with zinc is usually more severe than with aluminum, and corrosion protection coatings such as zinc or cadmium or other heavy metals not only can promote hydrogen uptake at crack tips but they enable co-deposition of hydrogen during deposition.
In studies on propagation of stress corrosion cracks in 18 percent nickel maraging steel Peterson and his associates found that when the steels was polarized to a potential of −0.77 V by coupling with a cadmium anode, the stress required for crack propagation in seawater was raised to a value close to that for crack propagation in air. However, when zinc at a potential of about −1.03 V was used as the source of the protective current, the stress value for crack propagation was reduced considerably below that for crack propagation in the absence of cathodic protection. This result suggests that caution should be exercised in applying cathodic protection in seawater to high strength steels of this type.
Moreover, UHSS components are usually coated for corrosion protection from environmental factors such as high humidity, changing ambient temperatures, salt spray, processing and operational chemicals. Sometimes these coatings may function as H permeation barriers. However, in operational service, coated components are seldom free of scratches that expose the bare steel to a galvanic couple potential. Sacrificial coatings cathodically polarize steel at the scratch producing hydrogen that could be absorbed at these sites. Low alloy UHSS suffer from hydrogen environment assisted cracking which can limit their use in marine environments.
FAILURE OF NICKEL BASE ALLOYS
There are at least eight nickel alloy systems of major commercial importance. These include pure nickel, nickel–copper alloys, nickel-chromium alloys, nickel-iron alloys, nickel–molybdenum alloys, nickel-chromium-molybdenum alloys, Ni-Cr-Mo-Fe-Cu alloys, and nickel based superalloys.36,37,38 Superalloys contain up to a dozen alloying elements. Many alloys were developed for high strength at high
37 K.G. Budinski and M.K. Budinski, Engineering Materials: Properties and Selection, 8th edition, Pearson, Upper Saddle River, N.J., 2005.
38 J. Kolts, “Environment Embrittlement of Nickel-Base Alloys,” p. 647 in ASM Metals Handbook, ASM International, Metals Park, Ohio, 1987.
temperature and low thermal expansion, but have also been adapted for seawater use due to good corrosion resistance. Some nickel base alloys are age hardenable while none respond to allotropic transformation type quench hardening. Nickel and copper are completely soluble in each other (Monels) and nickel exhibits good solubility with iron, chromium, manganese. Alloy 718 and Monel K-500 are two examples of alloys often used as high strength fasteners in seawater. Each alloy is based on a face-centered cubic austenitic matrix with good intrinsic corrosion resistance and often procured in the form of various studs and bolts in the solution annealed, cold worked and aged conditions.
Strength levels vary by alloy type and heat treatment but are 150 ksi (1034 MPa) yield strength and 180 ksi (1275 MPa) ultimate tensile strength with a Rockwell hardness of minimum RC 35-36 for aged 718. Cold work, annealing temperature and age hardening are crucial features of each alloy that affect hydrogen uptake, trapping, strength, and fracture resistance. For instance, the time-temperature-transformation curve for Alloy 718 illustrates how aging is performed after annealing. The annealing temperature is critical in that a grain boundary phase known as d, Ni3Nb, can be formed if annealed below a certain temperature. The alloy is aged hardened by formation of g’ and g’’ in the metallic matrix. Monel Alloy K-500 (UNS N05500) is a precipitation-hardenable nickel-copper alloy that combines the corrosion resistance of Monel Alloy 400 with greater strength and hardness. It is strengthened by a combination of cold work and precipitation age hardening to form fine coherent intermetallic Ni3(Ti,Al) phases.
General HE Susceptibility
Nickel-chromium-iron alloys containing more than about 40 percent nickel are not susceptible to stress corrosion cracking by concentrated chloride solutions such as magnesium chloride. A number of nickel base alloys have been observed to crack in boiling MgCl2, salt brines, seawater, HCl and H2S environments. These are reviewed in the ASM Metals Handbook, Volume 13, and the references therein.39
The principal type of EAC of nickel base alloys in seawater occurs when there is exposure to cathodic protection and hydrogen production. However, lab testing involving cathodic charging, the NACE solution, hydrogen gas, or other hydrogen producing environments can also cause susceptibility. The rule of thumb regarding the notion that hydrogen embrittlement in seawater requires a strength level greater than about 100 ksi (690 MPa) generally applies to nickel based alloys.
The hydrogen embrittlement assisted cracking (HEAC) behavior of Ni-based superalloys in gaseous H2 and aqueous exposure has been reviewed in detail. The potential severity of this cracking problem is apparent in nickel base alloys such
39ASM Metals Handbook, Volume 13, ASM International, Metals Park, Ohio, Table 16.
as Alloy 718.40 Subcritical environment-assisted cracking in high strength superalloys can be equivalent for gases and electrolytes as long as the dissolved hydrogen concentration is the same. This assumption is justified as the microscopic cracking modes were similar for each environment and involved a mixture of IG and TG slip-plane based cracking. However, occluded-crack chemistry analysis must be considered for exposure in seawater.
Figure K.6 shows that these thresholds are defined by a single function of the crack tip H concentration for alloy 718 stressed in aqueous-acidified chloride solution (•) as well as high pressure H2 (o). The KTH declines with increasing crack tip H concentration, above 20 ppm and independent of the crack tip environment that produced this H.
Similar data have been developed for solution heat treated and aged Monel K-500.41 The trend is the same but the details differ and must be understood to evaluate suitability in a given exposure environment. It should be noted that the crack growth rate slows considerably as the applied potential becomes less cathodic. In fact, at −0.7 V the crack growth rate is not detected in a lab test. Nickel-copper alloy bolt failures have also been observed numerous times in age hardened Monel Alloy K-500 (UNS N05500) subjected to normal cathodic polarization when coupled to aluminum anodes in seawater.42,43 Failures by intergranular cracking have been attributed to hydrogen production, uptake and embrittlement as a result of cathodic polarization and was first attributed to high thread root hardness of HRC 39 due to age hardening after threading. It was recommended that first annealing and then age hardening be conducted after threading to limit hardness below the acceptable hardness (HRC 35) of Monel Alloy K-500 (UNS N05500) recommended as a limit in sour systems. However, additional failures have occurred in roll-threaded Monel Alloy K-500 (UNS N05500) bolts annealed at 980-1050°C, water quenched, and precipitation age hardened at 500-600°C for 16 h producing a Rockwell C hardness of only 25. Embrittlement failures occurred in bolts after about one year under load to about 59 percent of the tensile yield strength and cathodically protected with anode grade aluminum. These failures have continued, and the root cause remains elusive.
40 L. Raymond, Fracture and Stress Corrosion Cracking Resistance of C465 (46 HRC), BioDur 108 (39 HRC), SpT 13-8 (37 HRC), SpT 13-8 (35 HRC), K-Monel 500 (31 HRC), and Zeron 100 (23 HRC), LRA Report #CTC’071024, L. Raymond and Associates, Newport Beach, Calif., 2008.
41 R.P. Gangloff, H.M. Ha, J.T. Burns, and J.R. Scully, Metallurgical and Materials Transactions A 45(10):3814-3834, 2012.
42 K.D. Efird, Failure of Monel® Ni-Cu-Al alloy K-500, Materials Performance 24:37-40, 1985.
43 L.H. Wolfe and M.W. Joosten, Failures of nickel/copper bolts in subsea application, SPE Production Engineering 3:382-386, 1988.
Dependence on Applied Potential
Hydrogen embrittlement under cathodic polarization in seawater follows a similar trend to steels and has been investigated for Alloy 718 (55Ni-20Fe-18Cr-5Nb-2(Al+Ti), (UNS N07718), and Monel K-500 (63Ni-27Cu-3Al-0.5Ti). An example of the combination of hydrogen concentration versus applied cathodic potential to predict Kth has been reported. The CH versus Eapp data was used to predict Kth and could also be adjusted to take into account the IR drop and alkalinity of cathodic crack tips leading to the family of curves relating hydrogen concentration to potential for various recess depths labeled as 0, 10, 5, and 100 mm in a rescaled recess with a 1 mm gap as shown in Figure K.7. Kth can then be predicted as a function of crack depth and applied potential in seawater for aged Alloy 718.
The predicted Kth within a crevice matches experimental data. Kth decreases as the cathodic potential decrease. The results for Monel K-500 show a similar trend and increased hydrogen production and uptake is seen at more cathodic potentials. Other high strength nickel base alloys are expected to display the same trends and is regulated by the details of the crack or occlude site geometry. This thinking can be applied to cracks and threaded geometries with confined spaces.