metal interactions with special reference to electrochemical approaches

Int. J. Hydrogen Energy, Vol. 6, pp. 119-138 Pergamon Press Ltd. 19~1. Printed in Great Britain © International Association for Hydrogen Energy

0360-3199181/030119-20 $02.00/0

HYDROGEN/METAL INTERACTIONS WITH SPECIAL REFERENCE TO ELECTROCHEMICAL APPROACHES H. J. FLIqT School of Physical Science, Hinders University of South Australia, Adelaide, Australia and J. O'M. BOCKRIS

Department of Chemistry, Texas A&M University, College Station, TX 77843, U.S.A.

(Receivedfor publication 16 July 1980) Ah~raet--A comprehensive review of hydrogen interactions in metals has been made with special emphasis on the electrochemical viewpoint. The object of this article is to provide the reader with a general knowledge of the physico-chemical aspects of hydrogen embrittlement of metals. Quoted papers can be consulted as an adjunct to this review, as all have been considered to be significant in advancing knowledge of hydrogen embrittlement.

1. I N T R O D U C T I O N THREE categories of hydrogen damage are classified as hydrogen embrittlement (Hirth and Johnson

[1]):

(a) Hydrogen environment embrittlement, owing to gas phase interaction of a metal with hydrogen. (b) Hydrogen stress cracking, which results from aqueous-solution-produced hydrogen. (c) Loss of tensile ductility, arising from dissolved hydrogen that tends to lower the fracture stress of the metal, but not resulting in a loss of yield strength. 2. DISSOLVED H Y D R O G E N

2.1. Crystallography of dissolved hydrogen 2.1.1. Neutron diffraction. Neutron diffraction has been used to study hydrogen in metals, and studies such as Bergsma and Goedkoep [2] gave indication of octahedral occupancy of hydrogen in the face-centred cubic lattice of palladium. Estimates of the Einstein frequency of hydrogen were made. Neutron spectroscopy is used for measurements of lattice vibrations of hydrogen/metal structures (Springer [3]). One can obtain from this evidence pertaining to the position of a hydrogen atom in a lattice, e.g. octahedral sites occupied in a face-centered cubic lattice (Oriani [4]). 2.1.2. Internal friction. Internal friction measurements (Gibala [5]) on hydrogen-charged iron produced a Shock relaxation peak at 45 K. Lord [6], by applying the linear solid model of relaxation peaks to describe the hydrogen Shock peak in a~-iron, has produced convincing evidence of interstitial hydrogen occupancy. Damping changes of the ratio, q~, are given by the equation: o)17 = A m 1 + 0)2r2' (1) where A m is the relaxation strength, ~ is the relaxation time and 0) is the angular frequency of the elastic wave. Observed values of ~)m~were (4-12) x 10-5, and this favors octahedral occupancy of hydrogen in oMron. 2.1.3. Partial molar volume. Even though hydrogen_atoms have a diameter of 1.06 A, giving a molar volume of 0.12 cm 3, the partial molar volumes, V, of hydrogen in metals are approx. 10 cm 3 g 119

120

H.J. FLITI" AND J. O'M. BOCKRIS

atom -1 (Table 1). The large value was first shown by Beck et al. [7], who found a value of 2.06 cm3g atom-1. 2.2. Summary of the state of dissolved hydrogen Dissolved hydrogen in metals occupies interstitial sites of the metal lattice, and these are octahedral sites in tr-iron. A large value of hydrogen partial molar volume shows that the dissolved hydrogen particles have a significant lattice expansion effect. TABLEI.

Solvent lattice

"~H (cm3g atom -1)

Pd flTi

1.73 4.1

/~Zn Ta ate

5.3 9.4 2.0

* Oriani and Bement [142].

3. ELECTRONIC STATE OF DISSOLVED HYDROGEN Isenberg [8] considered that dissolved hydrogen was a proton and used the free electron model of a metal theoretically to establish that no bound state exists for an electron in the resulting electric field of the proton. This is called the screened proton model. Valence electrons acting as a screen lower the energy of the system. Ebisuzaki and O'Keefe [9] have used this model to evaluate the solubility of hydrogen in transition metals and alloys.

3.1. Experimental evidence for a screened proton model 3.1.1. Magnetic susceptibility. Zanowick and Wallace [10] determined the magnetic susceptibility, X, as a function of increasing hydrogen content in vanadium. A temperature-independent X indicates that the hydrogenated metal exhibits paramagnetism. This depends on the density of states at the Fermi level, so that any donation of electrons to the conduction band by hydrogen should explain a change of X with hydrogen metal composition. Vanadium does give a decrease in the density of states, if the Fermi level is increased by accepting electrons. As the experimental magnetic susceptibility decreased with hydrogen content, it implies that the hydrogen atoms exist as protons. Jamieson and Manchester [11] have found difficulties in expressing X as a function of temperature for the Pd-H system when a rigid band model is assumed (Mott and Jones [12]). Wallace [13] considered that hybridization of ls orbitals of hydrogen to form states below the Fermi limit could explain the failure of the rigid band model. This has been theoretically tested by the augmented plane wave method (APW) by Switendick [14] to construct band structures for hydrides. 3.1.2. Nuclear magnetic resonance. Absorption of hydrogen by a metal as protons would affect the paramagnetism because of electron donation to the conduction band. Variations of the Knight shifts of group/IIB metals with absorbed hydrogen have been observed by Schreiber [15]. These show a decrease with increasing hydrogen concentration and can be interpreted as a lowering in the density of states in the conduction band by increased electron donation. Kazama and Fukai [16] observed that experimental Knight shifts measured for V-H, Nb-H and Ta-H alloys with high-resolution NMR are not explained by the rigid band model, but by means of bonding between hydrogen ls orbital and host metal 3d orbitals. Cotts [17] has found agreement of the rigid band model in explaining Knight shifts in Group VB metals with hydrogen.

H2/METAL INTERACTIONS WITH REFERENCE TO ELECTROCHEMICAL APPROACHES 121 3.1.3. Electrical resistivity. Initially, the solubility of hydrogen in palladium was limited by vacancies in the d band (maximum = 0.59 H/Pd atom ratio). Smith and Otterson [18] measured the resistivity of the palladium-hydrogen system past the 0.59 H/Pd atom ratio to the 1.0 H/Pd atom ratio. A model by Bambikidis et al. [19] to explain the resistivity of palladium as a function of concentration gave good agreement with Smith and Otterson's [18] results. Based on the electronic structure of Pd, the model gives a filling of the d bond of electrons from hydrogen atoms. The protons are screened strongly, as shown by measurements at temperatures in the range 42 K.

3.2. S u m m a r y o f the electronic state o f dissolved hydrogen Hydrogen in metals is protonic and donates electrons to the metal conduction band. The quantum mechanical formulation is not entirely consistent with this view (Kanamori [20]). Experimental findings indicate that a rigid band model is inadequate for describing the electronic situation. More work is needed.

4. HYDROGEN/HYDROGEN INTERACTIONS 4.1. Thermodynamics From Sievert's law: A ~ I ° - T/X-S ° = - R T I n c/p i,

(2)

where iron, A]~, the partial enthalpy, and A ~ , the partial entropy, are not functions of concentrations of gas pressure (Fast [21]). A classification of metals into exothermic and endothermic dissolvers of hydrogen has resulted, depending on sohibility/temp~rature characteristics. A treatment of solid/solution thermodynamics of hydrogen in metals has been given by McLellan [22] who treated a gas-metal-hydrogen equilibrium situation at constant volume and temperature. He applied it to various metals using the following expression for the chemical potential of hydrogen,/~, in solution (McLellan [23]): I~n = -EH -- TSn - k T l n kl -~/'

(3)

where En is the energy required to insert a proton into a lattice, with respect to the energy of the atom at rest in a vacuum, Sn is the partial excess entropy, and k T In [(0/fl)/(1 - 0//3)] is the configurational entropy, 0 being the ratio of solute to solvent atoms, fl being the number of interstitials for one lattice atom. Solubility data as a function of temperature for various metals were used to calculate En and St1- Excess entropies, Sn, were considered as vibrational, and their magnitude is affected by the dilation changes in the solute lattice; values of En are considered to be consistent with the proton model. McLellan [23] points out that the classification of hydrogen/metal solutions as endothermic or exothermic is misleading. En is always exothermic, but the energy of dissociation of a hydrogen gas molecule can be larger or smaller than EH. Some metal solutions with hydrogen are concentrated and long- or short-range interactions can occur between protons in the metal lattice. Lacher [24] used a statistical model of the solubility of hydrogen as a function of composition to describe the H-Pd system. By assuming the configurational entropy ideal and the partial energy of solution proportional to the concentration, he proposed an attractive (sic) interaction of the dissolved hydrogen particles. Simons and Flanagan [25] extended Lacher's [24] model assuming the proton model for absorbed hydrogen. They included an interaction term between neighboring protons, and a variable electron donation term. For the m-phase of Pd (Simons and Flanagan [26]), having a low hydrogen content, a negative deviation in Sievert's law was claimed to be explicable by an attractive interaction energy term, which the authors did not further rationalize. With the fl-phase of Pd, the electron donation term becomes more positive and cancels out the negative dependence of the attractive energy term. McQuillan [27] applied a statistical analysis to measured heats of solution, E, to explain the nature of hydrogen in bcc transition metals as a function of hydrogen/metal concentration. A plot

122

H. J. FLITT AND J. O'M. BOCKRIS

of E against composition exhibited a maximum, which was interpreted as due to the onset of significant repulsive interaction between protons. This was explained as an electronic repulsive force, and not as the effect of lattice distortion. An alternative model to that of proton/proton repulsion is given by Burch [28] who considered a strain energy term in the Simons and Flanagan model for the Pd-H system. Thus, this model relies on lattice distortion caused by the ingress of hydrogen and less upon the H + ~ H + repulsion. Good experimental agreement was found in the calculation of the Sievert's law constant, and the result is consistent with hydrogen having a large partial molar volume in Pd. Langeberg and McLellan [29] have calculated partial molar enthalpies and entropies for the Nb--, V- and T a - H solid solutions, and these can be explained by statistical blocking. A blocking model assumes that there is largely repulsioe interaction between hydrogen atoms in the metal lattice. A further study by Stafford and McLellan [30] has given more evidence that the predominant interaction must be repulsive, not attractive. McLellan and Harkins [31] have reviewed statistical models for higher solute hydrogen/metal solutions, in which more sophisticated models are considered. Models are based either on the positional entropy being ideal, whereupon the enthalpy describes non-ideality, or the enthalpy is independent of composition but the entropy is not. Wagner [32] has treated elastic interactions in hydrogen/metal systems and has shown that these may give rise to attractive interactions between protons. These are responsible for phase changes in metal/hydrogen systems. The situation remains somewhat contradictory. 5. H Y D R O G E N D I S L O C A T I O N INTERACTIONS 5.1. Internal friction Internal friction measurements by Gibala [5] have shown for iron that a cold work peak is present for hydrogen at low temperatures. This was explained as the relaxation of an interstitial, dislocation complex. By comparing the Snoek peak and the cold work peak, the equilibrium condition between hydrogen atoms as interstitials and those adsorbed at dislocations can be studied. The binding energy of hydrogen atoms to dislocations was determined to be 6.4 kcal tool -1. Sturges and Miodownik [33] found the value to vary between 4 and 6 kcal mo1-1. The variations arise because of the stress caused by atomic hydrogen in the metal (Miodownik and Sturges [34]). 5.2. Stress analysis The first analysis of stress effects on the solubility of hydrogen in iron was given by Beck et al. [7]. These authors established the relation: Co = Co exp Vo/RT,

(4)

where Co is the hydrogen concentration in the metal with no applied stress, Co is the concentration of hydrogen at regions of stress o (positive for tensile stress) in the metal, and V is the partial molar volume. Wagenblast and Wriedt [35], utilizing an X-ray method, measured V for hydrogen in iron as 2.2 cm 3 g atom -~, which agrees with the values calculable from the permeability data of Beck et al. [7] (Oriani [36]). Li et al. [37] formulated the thermodynamics of stressed bodies in terms of classical Gibb's thermodynamical equations. A more involved treatment for hydrogen introduced at constant overpotential was given by Bockris and Subramanyan [38]. Equation (4) was rederived to express Vn in terms of solubility of an unstressed site, Co, and a stressed site, Co, for hydrogen: R T a ln (Co~Co) = Vn, (5) ao where a is the tensile stress level in the metal. Conflict arose over the actual mode of deriving the formula, but the end result is the same (Chopra and Li [39], Bockris and Subramanyan [40]). Using a method for obtaining the permeability as a function of stress, Bockris et al. [41] obtained

H2/METAL INTERACTIONS WITH REFERENCE TO ELECTROCHEMICAL APPROACHES 123 results for the partial molar entropy for hydrogen in iron (-15 eV). The distribution of hydrogen around an edge dislocation was estimated from elasticity theory to show that, in the case considered, 20% of hydrogen associated with dislocations. Hirth [42] pointed out that the hydrostatic treatment as given above to dislocations should be modified, because hydrogen interstitials produce a tetragonal distortion field, image relaxations of dislocations are present in the metal, and this reduces the stress at the dislocation, and occupation of dislocation sites with hydrogen atoms is not a Boltzmann distribution, but a Fermi-Dirac distribution. These corrections lead to a considerable reduction in the amount of hydrogen associated with dislocations. 5.3. Mechanical properties A determination of dislocation/hydrogen interactions is given by testing the mechanical behavior of metals. The Portevin-le Chatelier effect for hydrogen-charged nickel (Wilcox and Smith [43]) results from hydrogen atoms interacting with dislocations and locking them. Blakemore [44] has considered an electrical driving force on hydrogen at dislocations in nickel during yielding. The interaction energy, U, is given by: U = 4/15E/q A

(6)

where Efis the Fermi energy of the metal, q is the charge on the screened proton, and A is the lattice dilation. Louthan et al. [45], in an assessment of tensile properties of metals with hydrogen, have concluded that dislocations play a major role in absorption of hydrogen on a localized basis, which is responsible for modifications of plastic deformation processes. As dislocations are mobile defects, Johnson and Hirth [46] have considered hydrogen delivery to voids or other traps by dislocation motion. They have argued that models for hydrogen pressure build up in voids and are not capable of explaining embrittlement at low hydrogen gas pressures. When applied to ferritic and austenic steels, the model predicts a small increase of hydrogen supersaturation at voids, but this was not considered sufficient to cause embrittlement. However, the authors appear to have neglected the effects of internal stresses upon the solubility of hydrogen. Tien et al. [47] have developed a dislocation transport model, considering Cotterell atmospheres of hydrogen at dislocations moving through the lattice and creating voids at inclusions. The model predicted the development of hydrogen pressures within voids, which would be rationalized with ductile fracture behavior. Experimental support for the earlier model has been found by Donovan [48], who observed hydrogen evolution from metals during plastic deformation. An alternative is that suggested by Subramanyan [49]. The stress levels at dislocations are regarded as sufficient to cause high local concentrations of hydrogen [Equation (4)]. These spread the voids and the joint of these is the mechanism of breakdown. 5.4 Summary of hydrogen~dislocation interactions Hydrogen/dislocation interactions occur. What is not resolved is whether the normal thermodynamic treatment of local solubility to local stress gives sufficient localized hydrogen to fill a void or whether it is necessary to supply more hydrogen by dislocation diffusion. 6. SURFACE/HYDROGEN INTERACTIONS In iron/carbon alloys, interracial surface areas are provided by carbide precipitates, and this can be extended to include matrix grain boundaries. Harhal et al [50] have explained exothermic occlusion of hydrogen in steels, for differing carbon content in terms of the internal surface adsorption of hydrogen at carbide/ferrite interfaces. The energetics of occlusion were found to be of the same order as those of chemisorbed hydrogen on iron surfaces. Hydrogen adsorption on the ferrite/carbide interfaces is sufficient to explain the observation without taking into account void formation.

124

H. J. FLIT1~ AND J. O'M. BOCKRIS

Newman and Shreir [51] studied the electrochemical entry of hydrogen into steels and explained the observed variation in permeation as the result of trapping hydrogen between ferrite and carbide particles. Gibala [52] has reviewed hydrogen/surface interactions, considering grain boundaries, stacking faults and twin boundaries as surface defects. Asaoka et al. [53] have used autoradiography to observe grain boundary/hydrogen interactions. Stacking faults are the major sources of trapping sites in ),-iron, and Miodowink [54] has reviewed hydrogen trapping at these defects. He suggests that hydrogen can: (a) reduce the energy of stacking faults in y-iron, (b) create hydrides, (c) produce deformation in martensites, and (d) affect the relative surface area of carbides. Craig [55] has suggested that in both martensite structures, hydrogen trapping is governed by elastic strain energy considerations at carbide precipitates, as annealing reduces trapping. This is consistent with a general loss of precipitate coherency. 6.1. Summary o f surface~hydrogen interactions

Surface/hydrogen interactions can be considered as major areas for further study. Compared to dislocation/hydrogen interactions, the amount of work done is less, but these defects may be as efficient in hydrogen trapping as are dislocations. 7. METAL VOIDS AND H Y D R O G E N That hydrogen gas exists in metal voids at pressures sufficient to cause the voids to spread is one of the older theories of hydrogen embrittlement (Zapffe and Sims [56], de Kazinczy [57]). Darken [58] has suggested that in acid-pickled mild steel, partitioning between voids and other defects occurs with hydrogen absorbed in the metal. Smialowski [59] has discussed the formation of hydrogen blisters under electrochemical charging conditions in steel and has concluded that stress is required for nucleation of voids. He agrees with Bockris and Subramanyan [60] that a critical fugacity is needed to create voids of significant size. The fugacity produced by the electrochemical overpotential has been considered to be in equilibrium with the equivalent pressure of molecular hydrogen within voids in the metal (Bockris and Subramanyan [60]). Corresponding to a critical overpotential is the critical pressure within voids to cause fissuring or rupture of the metal. Beck et al. [7] related anomalous permeation of hydrogen (Fig. 1, curve 3) to the formation of voids containing hydrogen under pressure. The critical pressure PH2 ~, was equivalent to the overpotential of charging at the onset of anomalous behavior (Fig. 1, curve 2). An expression for the expansion of a microcrack (void) was given by Bockris and Subramanyan [61] in terms of hydrogen pressure, PH2, as:

F. -2X - 1'

P.2 ~ L~(1 - : ) z J '

(7)

4O

::k

20 ~.xtxoX~x~ x~x (I) 0

4

¢, rain FIG. 1. Permeation-time curves, Armco iron, L = 0.77 mm, T = 25°C, 0.1 N H2SO4; (1) 4 × 10-4 A cm-2; (2) 8 x 10-4 A cm-2; (3) 4.3 mA cm-2 [Beck et al. [7]).

H2/METAL INTERACTIONS WITH REFERENCE TO ELECTROCHEMICAL APPROACHES 125 where Y is the Young's modulus, ? is the surface energy, l is half the crack length, and v is Poisson's ratio. Chene et al. [62], by using potentiostatic-controUed permeation of hydrogen in iron, found that embrittlement occurred at 232 mV of overpotential. This critical potential is related to a change in the electrochemical hydrogen evolution mechanism, as indicated by a change in the Tafel slope during hydrogen evolution (Fig. 2). Metallography confirmed the presence of voids only after the critical ooerpotential had been reached, in accord with the predictions of Beck et al. [7]. Dislocations are considered the sites of void formation as rapid straining of the metal produced no embrittlement. Thus, with sufficiently rapid straining, hydrogen has no time to segregate at dislocation sites or cause embrittlement. From Bockris and Subramanyan [60], the mechanism of hydrogen evolution can be pro- or nonembrittlement producing, due to the effective pressure (fugacity) corresponding to the overpotential of charging. However, the embrittling conditions arise not only as a result of a certain overpotential, but also because of the onset of certain mechanisms of the surface hydrogen evolution. Thus, in the work of Beck et al. [7], a change of condition which caused the permeation-time transients to change their character occurred at overpotentials more negative than about 250 mV, in reasonable agreement with the overpotential at which voids were found to begin. Roberts et al. [63] have described an in situ (Chene et al. [62]) tensile testing method with a scanning electron microscope, to observe void formation around inclusions in steel during plastic deformation. This method could be applied to study hydrogenated steels and would probably supplement the evidence of hydrogen void formation (cf. Chene et al. [62]). 7.1. Summary o f metal ooids and hydrogen

A void containing molecular hydrogen at high pressure has finally received direct evidence of existence, and such existence has been correlated with the electrochemical overpotential which has to reach a critical value for the voids to form and fill with hydrogen. The overpotential which creates such voids depends on the mechanism of the surface reaction. The in situ method of direct observation of void growth (Roberts et al. [63]) could be a possible way of giving additional evidence for the voids.

(1)/

o/o,,,o/

./

i0-(

d~) dtog P = - 417 mV

/

{IT) d~ jolI/ dLog p=- 98mV

%

I ~ io_5 ao

nO-4

I -I00

I -200 ~

1 -300

I -400

,

mV

FIO. 2. Permeation overpotential relation of iron (Chene et al. [62]).

126

H.J. FLIq'I" AND J. O'M. BOCKRIS 8. H Y D R O G E N DIFFUSION AND P E R M E A T I O N T H R O U G H METALS

8.1. Lattice diffusion

Hydrogen mobility in metals is of the same order as that of water molecules in the liquid phase, and so it is orders of magnitude above the mobilities of other interstitials. The activation energy for diffusion is small (2-5 kcal mol-'). Four temperature-dependent types of diffusion processes are possible (Kehr [64]), namely: (a) very low temperature band propagation by coherent tunneling; (b) thermally activated tunneling (incoherent) at low temperatures; (c) thermally activated jumps over a barrier at room temperatures; (d) fluid-like diffusion at high temperatures. Various methods are used to obtain diffusion characteristics of hydrogen in metals. SkOld [65] reviewed the use of quasi-elastic neutron scattering (QNS), which gives time and space development of the jump step in diffusion process (c), above. NMR spectra of hydrogen in metals can be analyzed theoretically to establish the mean residence time for a hydrogen atom on a lattice site (Cotts [17]). The most common method is to use a hydrogen concentration gradient across a metal membrane, either electrochemically (Bockris [66]) or by hydrogen gas pressure (Robertson [67]). The rate of hydrogen permeation through the metal membranes can be analyzed to give diffusion coefficients as well as the solubility as a function of fugacity (McBreen et al. [68]), which can be varied by more than 106 times by the use of simple electrical circuitry. 8.2. Hydrogen entry into metals 8.2.1. Electrochemical entry. Hydrogen is electrochemically evolved at metal interfaces by a number of reaction pathways (Boekris and Reddy [69]). All of these pathways involve adsorption of hydrogen atoms on the metal surface as a chemisorbed layer, a subsequent entry (Sakamoto and Miura [70]). McBreen and Genshaw [71] have given hydrogen evolution kinetic parameters for various evolution mechanisms for low (Langmuir) and high (Temkin) hydrogen atom coverages on the metal. The two main paths are: 2H30 ÷ + 2e~et~l~ 2Had~ + H20 (discharge), 2H~ds~ H2 (combination),

(8)

H3 O+ + e- ~ Haa~ (discharge), H 3 0 + + Had~ + e- ~ He (electrochemical desorption).

(9)

and

Bockris et al. [72] have found agreement with a coupled discharge mechanism on iron for the deduced permeation current kinetic factor. Kinetic factors for the permeation current corresponding to high and low coverages have also been determined (McBreen and Genshaw [71]). Bockris et al. [73] have considered the energetics involved in the hydrogen absorption reaction for various metals. At zero overpotential, the relation of the heat of permeation is: AHp = A H o + A H , ,

(10)

where A H o is the heat of diffusion and AH, is the heat of solution for the respective hydrogen metal system. Calculated heats of solution based on electrochemical permeation data agreed with values obtained from non electrochemical methods. 8.2.2. Gas phase entry. McCright [74] has expressed the gas phase entry of hydrogen at a rate, r:

r = uOM~ exp ( - A H * / R T ) ,

(11)

where u is a frequency term, 0 is the metal sites covered by hydrogen and Ms is the total number of metal sites covered. AH* is the heat of activation of the absorption process.

H2/METAL INTERACTIONS WITH REFERENCE TO ELECTROCHEMICAL APPROACHES 127 8.3. Electrochemical permeation of hydrogen 8.3.1. Diffusion. An experimental method of studying the diffusion characteristics of hydrogen in metals by electrochemical permeation was first experimentally developed by Devanathan and Stachurski [75]. Their electrochemical permeation cell is shown by Fig. 3. The transient response of hydrogen permeation, through metal membranes, can be used to obtain values of diffusion coefficients for hydrogen lattice diffusion. The diffusion coefficients can be evaluated from the time-lag to reach steady-state permeation (oxide film removal must first be assured) or the transientrise time constant for 63% of the steady-state permeation level. Using the time-lag method, the iron D-values appeared to vary with the thickness of the metal membranes. Wach and Miodownik [76] developed a method of correcting for the "barrier" effects. Wach [77] interpreted the "barrier" effect in terms of adsorbed hydrogen on the metal membrane (Ziichner [78]). 5

4

7

2 6 3 FIG. 3. Electrochemical permeation cell (Devanathan and Stachurski [75]).

McBreen et al. [68] developed an analysis of the permeation transients by solving Fick's second law, with defined boundary conditions for the metal membrane. Diffusion coefficients could be evaluated at any time of the permeating transient. Gileadi et al. [79] applied the above analysis to measure diffusion coefficients in platinum as a function of temperature (50-90°C). The major proportion of absorbed hydrogen was trapped in local strain regions, and only 5-15% diffused through the Pt lattice. Experimental improvements of the permeation technique have been made to obtain reliable diffusion coefficients, such as coating the input and output surfaces with palladium (Heidersbach et al. [80]). Kumnick and Johnson [81] used a Pd-coated electrochemical permeation technique to measure iron diffusion coefficients of deuterium and hydrogen in a temperature range of 9-73 °C. Early [82] has considered that with galvanostatic charging, the diffusion equation considered earlier for permeating measurements (MeBreen etal. [68]) was incorrect to assume an instantaneous hydrogen concentration at the metal surface. However, the time to achieve a steady overpotential can be less than milliseconds, so that no practical consequence arises from Early's [82] comment. 8.3.2. Trapping behaoior. The anomalous permeation transients observed by Beck et al. [7] were interpreted as being associated with hydrogen trapping, produced at a critical overpotential of hydrogen evolution. Bockris and Subramanyan [61] identified at least two distinct trapping sites by an analysis of the decay transient. The traps have different binding energies (Plusquellec [83]). An early theoretical treatment of trapping was derived by McNabb and Foster [84] who considered an equilibrium between hydrogen atom populations in lattice and trapped sites. Oriani [85] reformulated this treatment to take into consideration the effect of populations of hydrogen

H. J. FLIT/" AND J. O'M. BOCKRIS

128

atoms at traps ahflTattice sites. His result tended to show that dislocations and internal surfaces could account foPthe total amount of absorbed hydrogen in non-cold worked steel. Caskey and Pillinger [86] have solved McNabb and Foster's [84] basic equation of trapping to obtain a general analytical solution. The solution can be applied to permeation transients to obtain the nature of trapping. The authors thought that the time-lag estimates of diffusion coefficients and solubility may be inaccurate. Bockris and Subramanyan [61] considered a non-equilibrium situation with trapped and lattice hydrogen, because of the irreversibility introduced by the overpotential. The quantity of trapped hydrogen can be obtained by comparing the expected course of the permeation transient, and the observed anomalous transient at steady-state attainment. Figure 4 is an example of this situation, where the amount of hydrogen trapped can be calculated from the area contained between the tWO

curves.

E

0

I

I

I I

I

1

,

I

I

t.Ot(l in miffil.-',-~t(8 in rain-t)--~-

FIG. 4. A laboratory record of an anomalous permeation transient attaining a steady state. The broken line represents the expected round course of the transient in the absence of hydrogen trapping [5% Ni +95% Fe, ic = 105 mA cm 2, L = 0.094 cm, 70°C and 0.1M NaOH solution (Bockris and Subramanyan [61])]. The effect of temperature on the saturated hydrogen content and the diffusion coefficient has been studied by cathodic charging of steels by Newman and Shreir [87], using vacuum extraction. Results gave a decreasing saturation content of hydrogen with increase of temperature, which was consistent for trapping at surfaces or voids. High-temperature diffusion coefficients from hydrogen evolution curves during vacuum extraction were higher than those determined with electrochemical charging. The high-temperature diffusion coefficients were considered to reflect minimal trapping, and from this a simple model of trapping was developed. At a given temperature, the lattice hydrogen concentration was independent of steel composition, i.e. trapping occurs as molecular hydrogen. Marandet [88] has reviewed the effect of cold work on the permeation of hydrogen in steels. Cold working creates more trapping sites. McCright [74] reviewed the role of metallurgical variables, grain size and microstructure. The stronger the steel (re-alloying), the more soluble is the hydrogen, due to increased trapping. Pressouyre and Bernstein [89] have developed a kinetic model for hydrogen trapping in Fe-Ti and Fe-Ti-C alloys. Correlation to the time dependence of grain boundary cracking in these alloys can be achieved by the kinetic model during cathodic charging. Plateau behavior in the increase of grain boundary cracking with time is related to competition between reversible and irreversible traps demand for hydrogen. Hydrogen is taken from grain boundary traps to supply irreversible TiC traps, and so grain boundary cracking is arrested during this period. In relation to the permeation of hydrogen, this model is analogous to the prior work of Bockris and Subramanyan [61], who observed trapping for two distinct types of trap sites (Fig. 5). Pressouyre and Bernstein [90] have described various hydrogen trapping sites by means of permeation studies. Their work bears a relation to the discussion of trapping sites, for which experimental evidence was obtained by Bockris and Subramanyan [61]. 8.3.3. Electrochemical promoters of permeation. Bockris et al. [72] have suggested increased hydrogen permeation is created by CN, I- and napthalene because the M - H absorbed bond energy was lowered. There are three effects, as follows:

H2/METAL INTERACTIONS WITH REFERENCE TO ELECTROCHEMICAL APPROACHES 129 (a) An electrostatic interaction between the promoter and H- (chemisorbed). (b) Interaction of the promoter and the d-band of the metal (Conway and Bockris [91]). (c) The promoter adsorbs on sites favorable for hydrogen evolution (high bond energy) forcing evolution to occur at low-energy M - H adsorbed bond sites. Radhakrishan and Shreir [92] have studied Se, Te, S, Pb and Bi, which are listed here in decreasing effectiveness of promotion. The promotion effect is related to coverage effects, and the evolution kinetics on the electrode. Adsorption of the promoter will cause an increase of the M - H adsorbed bond due to restriction of sites on the metal surface (Parsons and Bockris [93]). Coverage increases on the metal surface, which will increase permeation.

40- t=O

o t I

I

I

I

I

=

16 min I

I

l (0.,'5 in rain-')

FIG. 5. De~y of permeation transient with trapping present (Bockds and Subramanyan [61]).

Smialowski and Szldarska-Smialowski [94] had postulated that hydrides are responsible for increased hydrogen entry. Newman and Shreir [51] have re-evaluated permeation promotion by the same species in steel. A sequence of promotion effectiveness should follow the bond strength of the hydrides formed. The saturation concentration of hydrogen with promoters showed a maximum as a function of pH. By plotting the maximum concentration as a function of promoter bond strength, an approximate linear relation was found. The hydride's role was explained as weakening the F e - H bond. McCright [74] lists various effects that must be considered important in the promotion mechanism, and these are as follows: (a) adsorption strength of species; (b) structure of the double layer; (c) surface and bond energy considerations; (d) character of hydride; and (e) electronic structure considerations. Kawashima et al. [95] have considered the promotion species for HES as undissociated molecular HES. The molecule adsorbed on steel acts as a bridge-forming ligand for proton discharge, which accelerates the discharge reaction, causing increased adsorption of hydrogen. 8.3.4. Inhibitors of hydrogen permeation. A series of organic nitriles was studied by Bockris et al. [72] to assess their inhibition of hydrogen permeation. As quantitatively similar effects are observed with aliphatic and aromatic nitriles, these molecules seem to adsorb vertically. Double layer thickening would result in hindering the approach of hydrated hydrogen ions to the metal surface. More energy thus is required to give hydrogen adsorption on the surface and would raise the overpotential of charging. Permeation decreases are explained by a reduction of the surface coverage. * Large molecules which can adsorb on metal surfaces are used as corrosion inhibitors. Singh and * This result does not agree with the later one of Bockris and Flitt [97]. Using a laser-pulse desorption technique, they found that the coveragewas largely unaffectedby the nitriles. Hydrogeninside the metal was greatly decreased.

130

H. J. FLITT AND J. O'M. BOCKRIS

Banerjee [96] have studied the effect of 2-mercaptobenzothiazole on hydrogen adsorption into mild steel. Absorbed hydrogen was measured using vacuum extraction. It decreased as a function of increasing inhibitor concentration. As the diffusivity, D, must be constant, the total hydrogen which diffused into unit area of the steel could be considered proportional toV~csat, where t is the time of charging and c~t is the metal hydrogen saturation concentration at the conditions of charging (Darken and Smith [98]). Hence, as a t i dependence of hydrogen adsorption is found with and without the inhibitor, the same electrochemical charging mechanism is thought to be present for these situations. Thus, coverage effect is considered the principal mechanism which causes inhibition of hydrogen absorption. Singh [99] has used a similar technique for the study of the inhibition of an Nsubstituted thiourea derivative, thiosemicarbarzide. The same t j dependence with the quantity of absorbed hydrogen, and a decrease in absorption with increasing inhibitor concentration, supports a coverage effect for the mechanism of the inhibition due to thiosemicarbarzide. Sherlock and Shreir [100, 101] investigated hydrogen entry into steel in phosphate solutions by varying pH, temperature and concentration of oxidizers for the phosphating process. The mechanism for hydrogen-reduced permeation is explained for the non-deposited phosphates in a manner similar to that proposed by Bockris and Potter [102] (effect of anion adsorption on M-H bond). 8.3.5. Summary of electrochemical permeation. Trapping characteristics of metals have been analyzed, and specific trapping sites characterized. Some promoters of hydrogen embrittlement have been shown to be related to a hydride film on the metal surface. Inhibitors can be considered to operate by a coverage effect which will increase the overpotential of charging, with loss of permeation due to the corresponding reduction of surface sites. 8.4. Gas phase permeation of hydrogen 8.4.1. Diffusion. Robertson [67] has studied gas phase permeation of hydrogen in nickel. The diffusion coefficient calculated was found to be independent of grain size, which varied from 30 to 120 tun. Gonzales [103] has retreated gas phase permeation data for iron. Not clarified is the degree of surface control over the permeation rates. Louthan and Derrick [104] have measured the permeability of deuterium as a function of temperature through austenitic stainless steels. With various surface preparations of metal membranes, the heat of activation for permeation, AHp, does remain constant. However, the permeability decreases with increasing surface finish. Oxide formation on the metal surface may play a role, Quick and Johnson [105] have compared the diffusivities of hydrogen and deuterium in iron from the gas phase. Classical behavior was evident at high temperatures, but not in the lower range. By Pd coating the input and output surfaces of a~iron, Miller et al. [106] were able to achieve diffusion-controlled permeation. 8.4.2. Trapping of hydrogen. Evidence of gas phase hydrogen-induced trapping in nickel has been found by Louthan et al. [107]. The effect of cold-rolling nickel was to produce dislocation sites which become effective traps, reversible in nature. 8.4.3. Promoters and inhibitors in the gas phase. Oxygen will inhibit embrittlement and, thus, could prove useful in practical applications (Liu [108]). Clermont et al. [109] have studied gas diffusion of hydrogen and the inhibiting effect of oxygen. The breakdown of the oxide film is suggested as the cause of delayed rupture. Thompson [110] has concluded that CO, CO2, N20 and SO2 are the most effective inhibitors of permeation. Acetylene is considered to be not an effective inhibitor, but may be useful in practice compared to other toxic ones. 8.4.4. Summary of gas phase permeation. Gas phase permeation has not been studied to the same extent as has electrochemical permeation. This area needs more research, as it is necessary to obtain safe gas transmission lines to establish a Hydrogen Economy.

H g M E T A L I N T E R A C T I O N S WITH R E F E R E N C E TO E L E C T R O C H E M I C A L A P P R O A C H E S 131

8.5.

Surface adsorption effects

8.5.1. Solution phase. Kim and Wilde [111] used a galvanostatic pulse technique to evaluate the coverage of adsorbed hydrogen on an iron surface: H adsorbed ~ H absorbed. Tilak et al. [112] have developed theoretical equations to represent overpotential decay behavior produced by anodic current pulsing on electrodes during hydrogen evolution. Relations have been established for the various electrochemical pathways possible for hydrogen evolution, and so information on the coverage of adsorbed intermediates is obtained. From Tafel behavior and the overpotential decay curves, control over the hydrogen evolution reaction by surface-adsorbed hydrogen can be found. Table 2, from Tilak et al.'s [112] paper, shows a comparison with Bockris and Subramanya'n's [60] estimation of fugacity. 8.5.2. Gas phase adsorption effects. Srikrishnan and Ficalora [113] have studied hydrogen adsorption from the gas phase, by measuring adsorption isotherms on iron. The adsorption process was reversible and gave a square-root pressure dependence, consistent with surface dissociation. Surface diffusion of anions of hydrogen was important to keep active adsorption sites free. Using magnetization chemisorption studies, Pecora and Ficalora [114] determined that the hydrogen adsorption is Temkin-like. Hydrogen bonding can be considered to be of two basic types, H + or H-, as revealed from paramagnetism. This agrees with earlier work that suggested hydrogen could be found to be adsorbed as two distinct species. The H- species is considered to be immobile, hence, this is interpreted as irreversible trapping of hydrogen on the metal surface. The authors give an explanation of hydrogen embrittlement in terms of irreversible H- traps forming within the metal. H ÷ species are considered to be mobile and supply the void with hydrogen. 8.5.3. Summary of adsorption of hydrogen. Few studies have attempted to look at adsorption in relation to hydrogen absorption (sic). Hydrogen entry is crucial to the embrittlement process. If ways of limiting adsorption of hydrogen on the surface were found, hydrogen embrittlement could be reduced. TABLE2.

Fugacity (~2) (atm) In terms of F

d

r/

0

Remarks

RTdlni Fast discharge-slow recombination

1/2(1 - 0)

{+2~IF~ exp ~

/

[+2~F~

Fast discharge-slow electrochemical desorption

1/(1.5 - 0)

Slow diseharge--fast recombination Slow discharge-fast electrochemical desorption Coupled discharge-electrochemical desorption

1/(0.5 + 0) 2 2

/+2r/*F~ exp ~ - - - ~ }

Coupled discharge-recombination

2 - 0 00) (1_1~

101.5exp +r/F

exp ~ - - ~ - ] 1 [-2r/F~ e~p ~ - - - ~ )

( O l~.0s~ 2 "1 0 Ok /

Nernstequation valid. Could be embrittling [ O._~l-OR~2 Nernst equation \ 1 - 0 OR / valid. Could be embrittling -Non-embrittling ( 0 1 -- _0~2 Non-embrittling 1 - 0

0R /

[ 0 1 - 0R~2 Only predictable if r/* \ 1 - 0 OR ] is known experimentally - 0

OR !

embrittling

kb = 0; g = 0; fl = 1/2; r/*--potential at which fast discharge-slow electrochemical desorption mechanism changes to coupled discharge-electrochemical desorption mechanism; OR-- coverage at the reversible potential; r/ is positive for cathodic processes.

132

H.J. FLITF AND J. O'M, BOCKRIS 9. MODELS OF H Y D R O G E N E M B R I T I ' L E M E N T

9.1. Pressure model

Zapffe and Sims [53] related hydrogen embrittlement of steel to the formation of high-pressure molecular hydrogen gas within the metal, which gave metal cracking and flaking. Theoretical calculations based on the equilibrium between atomic and molecular hydrogen gave sufficient pressures to fracture steel. Formation of molecular hydrogen was proposed to occur at defects, and the effects of aging and low-temperature annealing were considered the result of reversible trapping. A mathematical model of the pressure concept was developed by de Kazinczy [57], who considered hydrogen gas as expanding a Griffith crack. The fracture stress of the metal was lowered by the release energy of the expanding gas. Crack spreading was accompanied by hydrogen diffusion to the crack tip which explained temperature and time variations of embrittlement. Garofalo et al. [115] have used a Griffith crack, formed by dislocations, which spreads as a result of hydrogen gas pressure. The crack is found to reach instability at a critical pressure, P,~t, as follows: Pcrit = [" 2._G_yy ]1/2 [:r(1 - u)lcn,] '

(12)

where G is the shear modulus, u is Poisson's ratio, Y is the surface energy and It,, is the critical length for instability. Surface adsorption effects were also considered in the above analysis. This expression differs from Equation (7) by the alternate use of sheer stress, instead of tensile stress. Hence, the interchange of Young's modulus, Y, and shear modulus, G. Bilby and Hewitt [16] have repeated the Garofalo et al. [115] analysis of pressure spreading of cracks and have obtained a lower hydrogen fugacity needed for embrittlement than the previous authors found. Beck et al. [7], from electrochemical permeation findings, suggested that hydrogen produced from the metal corroding surface produces local high hydrogen concentrations in the metal just below the surface. This will result in blisters of molecular hydrogen forming and expanding to reach the surface and initiate a crack. The crack will then be subject to anodic dissolution at the tip (this being in turn subject to film formation and passivity effects), and hence, will enlarge. Tetelman [117] developed the thermodynamic pressure model to explain various aspects of crack growth. He predicted that cracks form and lengthen, until the Griffith criteria is reached, by internal hydrogen pressure in voids. Once at the Griffith criteria, crack growth is spontaneous, but will stop due to the loss of the hydrogen supply at the crack tip. Hence is explained the observed discontinuous cracking behavior (Fig. 6). Van Leeuwcn [118] has considered a model of hydrogen-induced grain boundary cracking by hydrogen. Voids are formed at grain boundary precipitates due to mechanical influences, such as the piling up of dislocations or grain boundary sliding. Hydrogen dissolved in the grain will tend to precipitate into the formed voids, and so a flux of hydrogen toward the void will occur. The

~Microcrock

FIO. 6. Microcrack nucleus formed by dislocation pile-up ahead of an advancing crack (Tetelman [117]).

H2/METAL INTERACTIONS WITH REFERENCE TO ELECTROCHEMICAL APPROACHES 133 effect of stress at such areas of the conflux of dislocations in influencing solubility is not taken into account. Using fracture mechanics, a model of hydrogen-induced grain boundary cracking is deduced. As the hydrogen coming to voids is contained in a limited volume of metal, adsorption occurs on the void walls, and the fugacity is higher than the mechanical pressure (Van Leeuwen [119]). No void embrittlement at grain boundaries is supported. The adsorption criteria in the model will tend to be eliminated as soon as complete coverage is obtained. The void pressure concept will be valid under this circumstance (Oriani and Josephic [120], Cialone and Asaro [121]). 9.2. Reduction of metal cohesive strength 9.2.1. Adsorption model. Petch and Stables [122] first suggested that reduction of the cohesive strength of iron with hydrogen resulted from surface adsorption. Petch [123] developed the model, using the Gibbs adsorption equation with a Langmuir isotherm. By equating fracture stress and grain size with the model, support for it was obtained by comparison with experimental results. Heady [124] has reworked the model given by Petch [123] to account for pressure in terms of fugacity, larger observed surface energies in iron, and a more updated adsorption isotherm. By relating the model to the stress intensity factor, an experimental evaluation of the model could be determined. The theory of Petch and Stables [122] is sound, but only accounts for a small part of hydrogen cracking in high strength steels. 9.2.2. Decohesion model. Troiano [125] suggested that embrittlement could occur by the cohesive strength of the metal lattice being lowered at regions of high stress. Areas of triaxial stress ahead of a crack tip would have a high hydrogen saturation, and thus provide a cleavage path for the crack to follow. Oriani [126] developed a mechanistic theory of lattice decohesion crack growth. Large concentrations of hydrogen are drawn to the crack tip owing to the elastic stress affecting the chemical potential of hydrogen, as shown by Beck et al. [7]. Velocity of crack growth is related to the accumulation of hydrogen by diffusion to the crack tip. Oriani [127] has compared decohesion to surface energy lowering, and concluded that the only difference is that the latter does not define a mechanistic path. Experimental verification of the model was obtained by Oriani and Josephic [128] who found that threshold hydrogen pressures and stress intensity factors were in agreement with the decohesion model. Van Leeuwen [129] has given a more rigorous treatment of Oriani's decohesion model by considering the reduction of the cohesive strength as a power function of hydrogen concentration and accounting for crack tip blunting by plasticity in terms of fracture mechanics. However, in this model, an imaginary elastic stress is considered to exist for a short time prior to plastic deformation giving ductile fracture, which is difficult to rationalize. When applied to Oriani and Josephic's [128] results, the model gives good agreement. Also, the prediction of Oriani's model that cracking will initiate at a finite pressure and zero stress intensity is eliminated, as the Van Leeuwen model needs infinite pressure to obtain this situation, i.e. stress is always needed (it may, of course, be local). 9.3. Dislocation models of embrittlement 9.3.1. Dislocation transport. Bastien and Azou [130] considered from tension tests on iron and mild steel that embrittlement was the result of hydrogen atoms segregating at Cotterell imperfections of dislocations, under the influence of metal plastic deformation. Deposition of hydrogen atoms in a cavity would create a void containing hydrogen gas, which under pressure would cause metal rupture. Segregated hydrogen could return to the lattice after a sufficient period of time, which would not result in void formation. Also, reversibility of embrittlement could be obtained by heat-treating the metal. Louthan et al. [45], in an assessment of hydrogen embrittlement, have considered that high hydrogen concentrations result from dislocation hydrogen interactions, which modify plastic deformation processes. Tien et al. [47], using the concepts of Bastien and Azou [130], have developed a dislocationtransport model of hydrogen embrittlement. Enrichment of plastically deformed regions of metals with hydrogen occurs whenever a dislocation loses its hydrogen atmosphere. This is possible if the

134

H. J. FLIT1~ AND J. O'M. BOCKRIS

dislocation velocity is sufficient to leave the hydrogen atmosphere behind, or the dislocation is annihilated, and when hydrogen atoms are transferred to a stationary sink, such as a precipitate or void. The last case is considered by the authors, in which voids are created which are irreversible and contribute to local fracture processes. Hydrogen gathers at inclusions, and the pressure builds up in the resulting voids. Using experimental results of strain rates in hydrogenated metals, the model is shown to be consistent with ductile fracture behavior. Johnson and Hirth [1] have considered a hydrogen transport model in which dislocations are annihilated. As hydrogen is lost by diffusion into the metal lattice, a balance is set up between hydrogen arrival and departure at void sites, with dislocation annihilation. However, because of the thermodynamically reversible nature of trapping assumed, the basis of the prediction is not very realistic. If reversible trapping behavior could be produced by alloy design, the model indicates that embrittlement would be reduced. 9.3.2. Hydrogen-assisted cracking. Beachem [131] has presented a model for hydrogen embrittlement, whereby concentrated hydrogen ahead of the crack tip aids whatever deformation process the microstructure will permit. Microvoid coalescence, quasicleavage or intergranular fracture are supported by deformation processes, depending on the stress intensity factor, and the concentration of hydrogen at the crack tip (Fig. 7). This model suggests that hydrogen does not lock dislocations in place, but unlocks them and allows multiplication and movement through the metal. Thus, hydrogen aids the mechanisms of extremely localized crack tip plasticity. Further evidence of the model from fractography has been given by Beachem [132].

(a)

(b)

(c]

(d)

Flo. 7. Sketches of microscopic fracture modes observed in these experiments as a function of decreasing stress intensity factor and concomitant decreasing cracking rate: (a) high K, MVC; (b) intermediate K, QC; (c) low K, IG; (d) IG cracking with an assist from hydrogen pressure. (Beacham [131].)

9.4. Relation to stress corrosion cracking Hanna et al. [133] first suggested that a possible mechanism for stress corrosion cracking could be in terms of hydrogen embrittlement. The obvious cathodic reaction in corrosion would be hydrogen evolution, and so large amounts of hydrogen could be absorbed. As discontinuous crack growth was observed, it was considered that corrosion-produced hydrogen was the source of embrittlement in aqueous environments. Van der Sluys [134] found that under anodic charging conditions, for A1S14340 steel, both hydrogen embrittlement and electrochemical corrosion appear to operate with the latter more dominant at high pH and anodic potentials. Brown et al. [135] measured the solution pH within a stress corrosion crack and found a value of 3.8 for steel, independent of the solution's external pH. Using this technique, Smith et al. [136]

HJMETAL INTERACTIONS WITH REFERENCE TO ELECTROCHEMICAL APPROACHES 135 studied the electrochemical conditions of a stress corrosion crack in AISI 4340 steel. The pH and potential measured at the crack tip are thermodynamically favorable for the evolution of hydrogen from water. Thus, it was considered that hydrogen embrittlement was always the mechanism of failure. Nielsen [137] has considered the role of hydrogen as a factor in stress corrosion cracking. He reported the observation of hydrogen gas from stress corrosion cracks in austenic stainless steels, and has pointed out that hydrogen evolution must be the primary cathodic reaction in this ease. Wilde [138] observed that the activation energy for crack growth of high-strength martensitic stainless steels is 9.5 kcal mo1-1 under anodic potentials which correspond to the value for hydrogeninduced crack growth under cathodic conditions. Conditions are always suitable for hydrogen evolution at crack regions, and, hence, the mechanism operating is considered to be hydrogen embrittlement (Flitt et al. [139]). Mehta and Burke [140] considered the role of hydrogen in stress corrosion cracking of austenic stainless steels by studying mechanical properties. They have rejected the concept of hydrogeninduced martensite. The crack tip is considered anodic with hydrogen evolved cathodically on crack walls, which ultimately diffuses to the crack tip. Flitt et al. [139] have suggested that observed discontinuous cracking in stress corrosion cracking is the result of void formation at the hydrogensaturated crack tip. The crack tip is considered anodic with the crack walls cathodic. Dissolution at the crack tip is a continuous process of crack growth, but on finding a void-saturated region, discontinuous crack growth can occur. Hydrogen is forced to the crack tip by the stressed region surrounding it, causing supersaturation in front of the crack tip and probably bond weakening theremhence, an increase in exchange current density (Scully [141]). These last models are also related, in their quantitative forms, to the content of the paper of Beck et al. [7], because hydrogen-assisted cracking models depend, essentially, upon the enhancement of the concentration of hydrogen near the crack tip, owing to the extra stress developed there. The concentration is not always explicit. Thus, the model of stress corrosion cracking is one of multi-determination. The original electrochemical model, in which the crack tip dissolves anodically and hydrogen is evolved on the crack sides, remains a central model. However, two other aspects must be introduced, both dependent on hydrogen. The first is that the simple electrochemical model provides cracks which move too slowly. It is only by assuming hydrogen has an effect on the tip regionmweakening FeFe bonds and increasing/0--that the orders of magnitude in velocity can be rationalized. Qualitatively, also, hydrogen diffusion to the crack tip has to be necessary to allow for the initiation period. The equation which governs these effects is: Co = coe ~"°/sr.

(13)

It is this which makes very high solubility effects possible in local regions. The equation retains validity over shorter times, even if u is greater than that for plastic deformation. However, there is certainly an additional non-electrochemical element in cracking (cf. cracking in insulators, such as glass). Thus, cracks make sudden forward movements, and this is described phenomenologically as tearing. It seems likely that this is hydrogen-dependent and that the crack moves ahead through the atmosphere of dissolved hydrogen at the crack tip by a non-electrochemical process (stress at the crack tip breaking Fe-Fe bonds). There is much which still has to be known about cracks; in particular, a verification as to the hydrogen content at the crack tip, but also work on constituents inside the crack. Acknowledgements--Thanks are due to Glen Hildreth and Lynn McCartney of Texas A&M University for their assistance in the preparation of this manuscript.

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