Hydrogen embrittlement (also referred to as hydrogen embrittlement cracking HEC, hydrogen stress cracking HSC, and hydrogen-induced cracking HIC) is a type of fracture that occurs in certain metals due to decreased toughness resulting from the absorption of hydrogen. The brittle failing of the material is due to diffusion processes (Fig. "Diffusion processes"), and requires sufficiently low strain speeds. Contrary to other types of embrittlement, notching loads are unsuitable for verification (Fig. "Verification of hydrogen embrittlement"). If brittle hydrides are created, as with titanium alloys, then the fracture behavior will be brittle in tests with high strain rates.
The influence of hydrogen can increase the notch sensitivity and crack growth rate considerably (Fig. "Diffusion processes"). In order to enter into the atomic lattice of a metal, the hydrogen must be atomic, since the molecule is already too large. Because hydrogen atoms immediately combine into molecules, the hydrogen atoms must form directly on the surface of the metal. This can occur in very different ways Fig. "Absorption scenarios"). The collection and recombining of hydrogen at flawed areas (such as non-metallic inclusions and inhomogeneities) in the structure creates a tensile stress concentration and causes the lattice to expand. The molecular hydrogen in the metal pores is under high overpressure and “bursts” the structure (Fig. "Effects of hydrogen on the structure"). The tendency to hydrogen embrittlement usually increases with the strength of the affected materials (Fig. "Sensitivity to hydrogen embrittlement").
Influence of materials:
Many technical materials are sensitive to hydrogen embrittlement to some degree, whereby the important influences for the action of the hydrogen can be very different. Al and Cu alloys are relatively resistant to hydrogen embrittlement, but in high-strength Al alloys, especially, embrittlement caused by hydrogen action on crack development has been observed. Austenite steels are affected to a certain degree. Ti alloys, ferrite steels, and Ni-based super alloys can embrittle strongly (Refs. 5.4.4-1 and 5.4.4-2). In high-strength steels, hydrogen embrittlement can cause failures at only 1 ppm hydrogen content above the solubility limit (for iron at room temperature, roughly 10-3 ppm), in Ti alloys, at about 35 ppm (solubility limit for titanium is 20-20 ppm). Aside from roughness, development of oxide films, absorption of hydrogen due to impurities, flaws in the micro-structure, and the hydrogen transport related to dislocation movements are all important influences on the hydrogen absorption (amount and distribution).
The greater the strength of a material, the greater its sensitivity to hydrogen embrittlement usually is (Fig. "Sensitivity to hydrogen embrittlement"). At the same time, it must be remembered that the alloy composition, as far as it influences strength, also affects the tendency to crack formation. For example, hydrogen embrittlement damage is especially pronounced in spring steels and maraging steels, whereby elastic properties and hardnesses are not influenced. Therefore, hydrogen embrittlement cannot be verified based on these characteristic values. The dynamic strength of steels at sufficiently high frequencies also seems to exclude any influence of hydrogen embrittlement (Ref. 5.4.4-13). With Ti alloys, however, there could be an influence if brittle hydrides are present on the surface and/or in the structure.
Fine-grained material is less sensitive to hydrogen embrittlement. This can be seen as due to higher ductility, lower density of “hydrogen traps” (larger sum of grain boundaries), and the more homogeneous stress distribution in the micro-zone.
The heat treatment state of the material is of special significance. For example, the tendency to hydrogen embrittlement decreases along with the fineness of spherical carbides in steels.
The considerably lower solubility of hydrogen in titanium materials at room temperature than at high temperatures or even in the molten state (welding, casting) means that at higher concentrations with slow cooling, the solution in the lattice and the enrichment at micro-defects leads to the precipitation of hydrides. Hydride development can be suppressed through rapid cooling. In some cases, the hydrogen embrittlement is related to the breaking-open of the hydride platelets (Ref. 5.4.4-3). Therefore, it is understandable that the size and orientation of the hydride platelets is important for the embrittlement effect. The limit values for hydrogen in titanium materials are: in alloys such as TiCu2, 100 ppm H; and in / alloys, 125 ppm H (Ref. 5.4.4-5). If hydride precipitations are detectable in metallographic cross-sections, then the hydrogen concentration is considerably greater than 100 ppm. Non-oxidizing acids on titanium surfaces create externally undetectable hydrides, which decompose at temperatures between 200°C and 300°C. The hydrogen that is freed during this process can enter into the metal. The absorption of hydrogen is blocked by oxide coatings. In titanium materials, suitable heat treatment in a vacuum can expel hydrogen (even recombined hydrogen). Relative to degasification (also called de-embrittlement) of steels, the required temperature is high (e.g. 700°C 6-24 h in a vacuum of 10-3 - 10-4 torr).
Unlike the behavior of steels (Fig. "Verification of hydrogen embrittlement") titanium alloys embrittle due to the development of brittle hydrides, and not due to deformation of the lattice. The hydrogen embrittlement of a-titanium alloys (such as TiCu2) is determined by a pronounced tendency to form hydride platelets in the structure. This also causes pronounced brittle behavior during fracture even at high strain rates. Both a decrease in notch impact strength and decreased plastic strain in tension tests have been observed. Therefore, verification of hydrogen embrittlement in -titanium alloys is done with notch impact tests. Time-dependent failure under steady stress is not pronounced.
Even * and /- titanium alloys (such as Ti6Al4V) can form hydrides and embrittle due to hydrogen. However, with these alloys the hydride development is not as pronounced as in a alloys. In this case, as with steels, dissolved diffusible hydrogen in the lattice seems to be more embrittling than the hydride development. The hydrogen that influences the crack development forms primarily on the boundary surfaces of the /-phase. Hydrogen diffuses at the crack tip, where brittle carbides form and promote crack growth. This makes understandable the effect whereby decreasing expansion rates promote crack behavior. Depending on the strength of the affected material, the hydrogen content, and the sharpness of any present notches, there is a static fatigue limit below which no fracture will occur, regardless of the temporal duration of the loads. Verification of hydrogen embrittlement in /-titanium alloys is done with notching rig tests.
Influences of surrounding and operating factors:
Hydrogen embrittlement is especially pronounced at “medium” temperatures (Ref. 5.4.4-3). For example, hydrogen embrittlement is only observed in ferritic steels in the temperature region between -70°C and +140°C (Ref. 5.4.4-4). This can be explained by the temperature-dependant mobility of hydrogen. The following influences promote the effect of hydrogen on crack development in metals:
The greater the surface roughness, the greater the surface area available for the inward diffusion of hydrogen. Therefore, rough surfaces also increase the tendency to hydrogen absorption and crack initiation. Hardened surfaces (e.g. through shot peening), on the other hand, generally have compressive residual stresses with metal lattice tightening that slows the inward diffusion of hydrogen. If, during hydrogen absorption, the surface relatively quickly precipitates largely hydrogen-proof coatings, it can considerably reduce the amount of inward-diffused hydrogen, relative to cases with slow or porous coating development. This can occur in identical processes merely by changing the precipitation parameters, for example.
Special attention should be paid to surface contaminants, if these are to be removed in baths. In one case, iron oxide deposits on compressor blades made from a high-strength titanium alloy embrittled due to hydride formation in an HNO3 bath (Fig. "Surface contaminants"). This is especially surprising since no major hydrogen formation was expected to occur in this bath.
Figure "Diffusion processes" (Ref. 5.4.4-3): Hydrogen that has been absorbed by metals can considerably increase the growth rate of cracks (e.g. also corrosion cracks, see Fig. "Role of hydrogen in corrosion processes"). Experts worldwide do not agree in their estimations of the role of hydrogen on the embrittlement process. To the author, one, albeit old, process seems to describe the damage process in steels so plausibly, that it will be explained in the following. It must not be denied that at least the explanation of the pore development phase is not considered to be proven by much technical literature (Ref. .5.4.4-4).
At the crack tip, a typically shaped plastic deformation zone forms (top left diagram), in which a lattice expansion occurs under the influence of the 3-axis tensile stress. Atomic hydrogen diffuses into this area. The hydrogen diffusion occurs (top right diagram, Refs. 5.4.4-3 and 5.4.4-6) in the direction of the greatest tensile stress conditions (largest lattice expansion) present in a “germinal point”. These are flaws in the transition of the plastic zone to the elastic zone. The hydrogen collects at the relatively small “germination point”. In less hard steels with hardnesses below about 30 HRC, the hydrogen can recombine to molecules that can no longer diffuse into the metal. In this way, hydrogen pressure pn (104 - 108 bar) rapidly builds up. When the gas pressure in the expanded (toward the inner crack) pores around the “germination point” in front of the crack has risen sufficiently, the “explosive effect” causes it to combine with the crack tip and increase crack growth.
However, this damage model is unsatisfactory for explaining the behavior of high-strength and case-hardened steels. For example, in these materials, damage has been observed to decrease again after longer times of conditioning (Ref. 5.4.4-14). The hydrogen is trapped at lattice deformations in reversible or irreversible “Traps”. This model also explains why long times of conditioning diminish the embrittlement effect, since the hydrogen is able to distribute better into the lattice from the reversible traps. However, the reversible traps can also act as hydrogen sources and provide hydrogen to zones with lattice expansions (e.g. in front of a crack tip). This explains the grain boundary cracks typical for hydrogen embrittlement.
The tensile stresses which cause hydrogen embrittlement must not necessarily be caused by external loads. Internal stresses from heat treatments, welds, coatings, and machining processes are frequently sufficient to cause hydrogen to collect in resting parts, such as during storage. Once hydrogen has collected at the “germination points” and created micro-cracks/pores, the damage will be irreversible.
As long as the hydrogen in reversible traps is atomically dissolved in the lattice, the damage is reversible (bottom diagram); suitable heat treatment (expulsion of hydrogen - degasification) can prevent the embrittlement process (Fig. "Disembrittlement").
Figure "Effects of hydrogen on the structure": The crack-promoting effect of hydrogen on metals can be caused by different effects (top diagram) that can occur individually or in combination with one another:
Deformation of the metal lattice: This embrittlement mechanism is dominant in high-strength and case-hardened steels (Ref. 5.4.4-14). The metal lattice is deformed by included hydrogen atoms. This reduces the mobility of dislocations, resulting in a loss of ductility. The hydrogen is caught in reversible or irreversible traps, depending on whether or not diffusible oxygen can be released. These traps are caused by lattice inhomogeneities and are located primarily in the grain boundary zone. This explains the intercrystalline crack development typical for hydrogen embrittlement. The hydrogen diffuses primarily in already-deformed atomic lattices. These include, for example, areas that have been cold-deformed under mechanical tensile loading, as is the case at crack tips (Fig. "Diffusion processes"). But even hardened martensite steels have suitable deformed lattices. These influences are especially effective in high-strength steels.
Recombination and gas pressure buildup: The hydrogen can diffuse out of the lattice into flawed zones (e.g. cast porosity, micro-cracks, deformation pores, inclusions, Ref. 5.4.4-6), around which the lattice is expanded due to the stress concentration, e.g. under internal stresses. In less high-strength steels (with hardnesses below about 30 HRC), pore formation is explained by a recombination to molecular hydrogen. This hydrogen gas can no longer escape through diffusion due to its relatively large molecular diameter. In the present or forming hollow spaces, extremely high pressures (top right diagram, Ref. 5.4.4-4) build up with an “explosive effect” which can cause brittle fractures without any major external loads.
A typical phenomenon is the microscopically small pores on the grain boundaries (with no sign of ductility of the crack honeycombs) of the fractured surface (bottom left diagram), which is REM-detectable evidence that hydrogen influenced the damage mechanism (Ref. 5.4.4-4). Hydrogen embrittlement is usually intercrystalline (bottom right diagram, Ref. 5.4.4-7). However, there are also cases of transcrystalline fractures. Hairlines (so-called crow`s feet) are flat crack honeycombs due to the remaining ductility of the material.
Hydride formation: Hydrogen can be chemically bonded by metals as a brittle hydride. Hydrides in the form of surface coatings or platelets in the structure promote brittle fractures. This embrittlement form is especially well known in Ti alloys (see).
Figure "Absorption scenarios": Atomic hydrogen can develop on metal surfaces and be absorbed by diffusion under various different scenarios:
Metal melts (“1”): At the high temperatures, the steam is split and atomic hydrogen develops at the surface of the melt. Especially large amounts of hydrogen can dissolve in melts, which leads to pronounced crack development in the frozen material.
Welding (“1a”): Contact between water and the melt is caused by
Casting (“1b”): In this case, also, hydrogen absorption occurs due to water from a damp environment. The causes are similar to those in welding. For example, as with damp electrode coverings, the use of damp shell moulds can provide a hydrogen supply.
Processing baths (“2”): In galvanic baths (e.g. coating baths for silver, chrome, nickel, cadmium, or zinc) or chemical baths (e.g. etching baths, electroless coating baths for nickel, etc.) with cathodic hydrogen creation at the part surface (cathodic reduction, H2 → 2 H+ + 2 e). These reactions also occur during certain corrosion processes (Fig. "Role of hydrogen in corrosion processes"). These include corrosive processes, especially at crack tips, where there is increased possibility for diffusion of the atomic hydrogen due to lattice expansion (Fig. "Diffusion processes") (elastic deformations, ductile cold strain).
“Contact diffusion” (“3”): Hydrogen diffusion can occur in case of intense metallic contact with a hydrogen-charged surface, such as screwed or shrink-fit part surfaces.
Thermal decomposition of hydrogen compounds (“4”): If a compound containing hydrogen, such as water, is heated to high temperatures, it may thermally decompose and result in the creation of atomic hydrogen. This hydrogen can diffuse into the metal walls of, for example, pipes or boilers.
During case-hardening through carbo-nitriding (Fig. "Gear failure"), the gas decomposes, creating atomic hydrogen (e.g. 3 Fe + CH4 → Fe3C + 4 H) that diffuses into the part being carburized.
Diffusion of catalytically dissociated hydrogen (“5”): From the molecular state (e.g. natural gas or hydrogen as engine fuel), dissociation due to a catalytic effect can occur at the surface of a part (e.g. a pipe). The dissociation can be explained by plastic deformations in the micro-zone, in which active fresh metal surfaces are created. These plastic deformations can be expected, for example, under LCF loads.
Figure "Role of hydrogen in corrosion processes" (Ref. 5.4.4-8): Hydrogen embrittlement plays a deciding role in the damage mechanism SCC (Chapter 126.96.36.199). During the anodic partial reaction, the actual erosive/dissolving corrosion process of the metallic material is accompanied by reaction coating development (Fig. "Coating damage I"). This creates positively charged metal ions when electrons separate (oxidation process).
In the case of anodic SCC the material is in a passive state, in which dissolving takes place only at the crack tip. At the surface of the part and crack, the negative charge of the material leads to an anodic partial reaction and atomic hydrogen creation due to the charge equilibrium with the positively charged hydrogen ions of the surrounding electrolyte (reduction process). This hydrogen can diffuse into the part and accelerate crack growth at the crack tip (Fig. "Diffusion processes"). The anodic and cathodic partial reactions area always coupled. In corrosion, the cathodic partial reaction usually determines the speed.
In cathodic SCC the material dissolution at the part surface occurs in zones with active corrosion behavior (anodic zones). The dissolving of metal and creation of positively charged metal ions frees electrons, which move to cathodic (passive re. corrosion) surface zones where charge balancing with the hydrogen ions of the electrolyte creates atomic hydrogen, which diffuses into the metal lattice and promotes SCC.
The creation of passive reaction coatings during the corrosion process can result in the development of protective conversion coatings (passive coatings). This allows the same metal to occur in both the passive and active state, depending on the corrosion conditions.
The influence of hydrogen embrittlement components in SCC can be seen on the fracture surfaces through characteristics such as micro-pore formation, quasi crack surfaces, and “rooster feet” (Fig. "Effects of hydrogen on the structure", Ref. 5.4.4-4). The passive state retards corrosion and is therefore usually desirable, but it is also more prone to deep localized corrosive attack, the dreaded pitting corrosion (Fig. "Tensile stress").
In Ti alloys, etching and signs of hydrogen embrittlement were discovered in connection with the influence of typical aviation-use hydraulic oil above 130°C (Ref. 5.4.4-13). At these temperatures, the hydraulic fluid evidently forms an acid that corrodes titanium and causes hydrogen embrittlement. According to the literature, only a certain beta-titanium alloy is immune to this damage. This alloy is used on the tail cone and thrust jet of the newest cargo aircraft.
Figure "Sensitivity to hydrogen embrittlement": Generally, the greater the strength of a material, the greater its sensitivity to hydrogen embrittlement (top diagram, Ref. 5.4.4-9). With increasing hardness, only a fraction of the amount of hydrogen that is diffused into the material is necessary to have a pronounced effect in reducing the contraction at fracture.
The bottom diagram shows the decrease in strength and life span in notched samples of a high-tempered steel under static loads. The samples were charged with hydrogen and subjected to varying degasification times at a heating temperature of 150°C. The influence of the heating times on the usable strength can be clearly seen.
Figure "Verification of hydrogen embrittlement": The effect of hydrogen on crack development in titanium alloys depends on the creation of brittle hydrides (see). If relatively slow diffusion processes occur in metal (e.g. steels), embrittlement (expansion of the lattice) is delayed (Fig. "Diffusion processes"). Therefore, hydrogen embrittlement is not usually (except for embrittlement of a-titanium through hydride formation) verifiable through notch impact tests or tensile tests (usually at high speeds). The high loading speed does not allow sufficient time for embrittling diffusion processes.
A decrease in crack strain and contraction at fracture can be expected from strain-controlled tensile tests with <0.05 mm/mm.min.
This damage in a creep test can be recognized by the decrease in strength and operating life.
The tensile notch test does not give any clear results.
On the other hand, influence on the crack development (critical loads for crack initiation and crack growth rates) can be expected from fracture mechanical CT samples with typical high stress concentrations (Fig. "Tensile stress").
Figure "Disembrittlement" (Ref. 5.4.4-5): As long as no recombination of atomic hydrogen to molecules has occurred (Fig. "Diffusion processes"), several hours of heating in air at a few hundred °C can expel the hydrogen or distribute it in the lattice in such a way that it sufficiently prevents the crack development effect. It is also possible to “disembrittle” Ti alloys that have become embrittled through hydride formation (see) by pyrolizing the hydrides at high temperatures (about 200° - 300°C). Because of the danger of recombination and/or bonding in irreversible “traps”, disembrittlement (see Fig. "Effects of hydrogen on the structure") must be done immediately (within a few hours) after hydrogen absorption in order to be successful.
Experience has shown that this is where the most mistakes are made, when the parts are allowed to sit for too long before disembrittlement or degasification is forgotten entirely.
The diagrams show that with increasing degasification temperatures and longer hold times, considerable reduction of the original hydrogen content occurs. The remaining amount of dissolved hydrogen depends largely on the amount of dissolved hydrogen before the disembrittlement process. There is no great difference between the degasification behaviors of chromed and cadmium-plated parts. However, it is obvious that cadmium-plated steels are less able to expel hydrogen than are hard-chromed parts.The difference may lie in the gas permeability of the coatings with a different sealing effect.
Hydrogen embrittlement can become noticeable in several different ways:
Typical damages in aircraft engines are:
Figure "Crack growth in strong materials": Typical parts in which hydrogen embrittlement damage is frequently reported are hardened or high-tempered parts made from low- and high-alloy steels with galvanic coatings, especially cadmium coatings (Ref. 5.4.4-15). The supply of atomic hydrogen seems to be especially intense during this process. Typical parts include disk springs, thread inserts, spring washers, plain washers, and threaded rods (top diagrams).
The lower diagram shows a hydraulic piston from a thrust jet adjuster made from a high-strength heat-treated steel. After a restorative chrome-plating procedure, certain areas (arrows) that were not chrome-plated developed cracks with typical characteristics of hydrogen embrittlement (Fig. "Effects of hydrogen on the structure").
Figure "Fracture of compressor disk": This integrally cast compressor disk from a helicopter turbine made precipitation-hardened high-strength steel (type 17 4 PH) was in an unfavorable heat-treatment state with too great a hardness. Hydrogen absorption probably occurred in an etching bath during new part manufacture. The disk failure occurred spontaneously during operation, and the disk broke into many pieces (compare with Fig. "Material behavior"). The crack development was transcrystalline with typical signs of hydrogen embrittlement such as micro-pores on material inhomogeneities.
Figure "Gear failure": Teeth fractures on case-hardened gears due to inner crack development following hydrogen embrittlement. The gears evidently absorb hydrogen during case-hardening (see Fig. "Absorption scenarios"). The damage then seems to occur as follows (Ref. 5.4.4-11): the high internal stresses in the case-hardened layer at the tooth surface induce high tension residual stresses in the tooth core (top left diagram). The smaller and more filigreed the tooth, the greater the danger that a circular hydrogen embrittlement crack will develop at internal notches within several hours. An intercrystalline hydrogen-induced brittle fracture then originates in this crack, and can later lead to a dynamic fatigue fracture under the operating loads, destroying the gearing.
Figure "Weld seams and turbine housings": Older engine types often have compressor housings that are welded constructions made from low-alloy heat-treated steels. Similar housings are found in modern engines near the outlet of the low-pressure turbine. The complex shape of these housings causes them to develop high tension residual stresses that are induced by welding during new part manufacture or repair procedures. If moisture (e.g. in the shielding gas or as condensation water) is present during the welding process (see Fig. "Absorption scenarios"), then hydrogen from the moisture can be absorbed by the melt (Ref. 5.4.4-12). This can lead to delayed crack initiation in the weld seam area during storage, manufacturing processes, or operation (arrows show typical crack-prone part zones). For this reason, if this type of crack is suspected, REM inspections of the fracture surface should be conducted first of all, in order to confirm characteristic signs of hydrogen embrittlement (see Fig. "Effects of hydrogen on the structure").
Figure "Surface contaminants": It has been discovered that, under certain circumstances, hydride formation can cause massive hydrogen embrittlement in high-strength titanium alloys even in oxidizing etching baths (such as HNO3). Important influences on this include:
In light of this, there is a question as to dangerous hydrogen embrittlement on titanium part surfaces in cleaning baths. Steel shot peening or machining with steel tools (e.g. polishing brush) can leave iron wear products/iron oxide on the part surface.
However, no such cases have been reported yet.
There are a number of preventive and corrective measures available to prevent damage due to hydrogen embrittlement:
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5.4.4-2 J.M. Bernstein, A.W. Thompson, “Effect of metallurgical variables on environmental fracture of steels”, periodical “International Metals Reviews”, December 1976, pages 269 to 287.
5.4.4-3 Tetelman, Mc Evily; “Bruchverhalten technischer Werkstoffe”, Stahl und Eisen Publishers mbH, Düsseldorf, 1971, pages 414 to 433.
5.4.4-4 L.Engel, H.Klingele, “Beitrag des Rasterelektronenmikroskops zur Beurteilung wasserstoffinduzierter Brüche”, periodical “Archiv Eisenhüttenwesen” 48 (1977) Nr. 10 (October), pages 555 to 560.
5.4.4-5 H.Simon, “Oberflächenreaktionen an Titanwerkstoffen”, periodical “Metalloberfläche” 4-1982, pages 211-217.
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5.4.4-7 V. Jürgens, “Beitrag zur Deutung der Rißentstehung bei verzögerten Brüchen im Schweißgut von Verbindungsschweißungen höherfester Stähle”, Dr.-Eng. dissertation, Technical University Braunschweig, 1975.
5.4.4-8 B.F. Brown, NRL Report 6041, Washington D.C. (1063).
5.4.4-9 M.R. Louthan Jr., “Role of Hydrogen in Stress Corrosion Cracking”, “Review in Coatings and Corrosion”, 2/3 (1977), pages 103 to 119.
5.4.4-10 G.E. Lukas,“Hydrogen Embrittlement in Medium and High Carbon Steels”, periodical “Finishing Industries”, Jan. 1977, pages 50-52.
5.4.4-11 R.Weiner, H.Fuchs, “Untersuchung zur Wasserstoffversprödung in galvanischen Bädern”, periodical “Metalloberfläche”, 25. 1971, Volume 1
5.4.4-12 T. Günther, H.Gräfen, “Wasserstoffversprödung von Feinkornbaustählen in Abhängigkeit von der Legierungszusammensetzung, der Gefügeausbildung und der mechanischen Belastung”, periodical “Werkstofftechnik”, 10, (1979) pages 373 to 390.
5.4.4-13 M.Peters, C. Leyens, J. Kumpfert, “Titan und Titanlegierungen!”, DGM, Informationsgesellschaft publishers.
5.4.4-14 S.Beyer, “Wasserstoffversprödungsanfälligkeit hochfester Werkstoffzustände am Beispiel einsatzgehärteter und angelassener Schrauben”, Darmstadt, Technical University dissertation 1997, Shaker Publishing GmbH, ISBN 3-8265-2917-0, pages 1-202.
5.4.4-15 “ASM Handbook”, Formerly Ninth Edition, Metals Handbook, Volume 13, Corrosion, pages 1037 to 1090.