The effects of irradiation on the microstructure and mechanical properties of nickel-based alloys are com­plex and, although the main factors affecting their behavior have been identified, a full understanding of radiation-induced effects remains elusive. This is par­ticularly true of the precipitation-hardened alloys, typified by Nimonic PE16 and Inconel 706, where the role of the hardening phases — which confer high ther­mal creep strength, but are redistributed during irradi­ation and may possibly influence swelling behavior and contribute to intergranular embrittlement — is unclear. The radiation-induced effects considered in this chapter-void swelling, irradiation creep, the evolution ofprecipitate and dislocation structures, and irradiation embrittlement — are interrelated in several ways, but particularly through the effects of point defect fluxes and the consequent redistribution of solute atoms.

The beneficial effect of nickel on the swelling resistance of austenitic alloys is well known, but a

clear explanation for the minimum in swelling that is found in alloys containing about 40-45% Ni has not been forthcoming. There is general agreement that the major influence of alloy composition on swelling arises through its effects on the effective vacancy diffusivity and on segregation via the inverse Kirkendall effect. However, on what appears to be the mistaken assumption that the swelling resis­tance of nickel-based alloys derives from an extended void nucleation period, swelling models have largely focused on factors affecting the nucleation rather than the growth of voids. Data for neutron-irradiated Nimo — nic PE16, for example, indicate that its swelling resis­tance is due to a combination of a comparatively low saturation void concentration, which is reached at a relatively low displacement dose, and a low void growth rate. The minimum critical void radius concept appears to offer the most plausible explanation for the minimum in swelling found at intermediate nickel contents, although experimental data comparing the behavior of PE16 and a nonprecipitation hardenable alloy with a similar matrix composition indicate that, in addition to the Ni content of the alloy, the presence of Si and/or the y’ forming solutes Al plus Ti may also be important. The dependence of the void growth rate on Ni may be related to the effects of radiation-induced segregation on the bias terms for the absorption of point defects at sinks, though again there is evidence that minor solutes, including Si, B, and Mo, as well as the y’-forming elements, have a beneficial effect on the overall swelling behavior of nickel-based alloys.

The irradiation creep behavior of nickel-based alloys generally appears to be similar to that of austenitic steels, though the higher thermal creep
strength of precipitation-hardened alloys is an advantage at higher operating temperatures. The main drawback of precipitation-hardened nickel — based alloys for reactor core applications is per­ceived to be a high susceptibility to irradiation embrittlement. Although it has been suggested that a combination of matrix hardening and of grain boundary weakening due to the formation of brittle intergranular layers of g0 (e. g., in the case of Nimonic PE16) or Z phase (in the case of Inconel 706) is responsible for the irradiation embrittlement of these alloys, there is strong evidence, at least for g0- hardened materials, that helium is the primary cause of low ductility failures. Experimental data have shown that the implantation or generation of rela­tively small amounts of helium can give rise to low tensile ductility, with intergranular failures initiated by either the growth and linkage of cavities or by wedge cracking depending on test conditions and helium distribution, under conditions where grain boundary g0 layers are not formed. However, irre­spective of the details of the embrittlement mecha­nism, it is evident that this aspect of radiation damage does not preclude the in-core application of nickel-based alloys, as has clearly been demon­strated by the successful use ofPE16 for fuel element cladding irradiated to high burn-ups in PFR. The long-term integrity of PE16 cladding is attributable to the dimensional stability of the alloy, arising from a combination of good swelling resistance and high creep strength, and relatively low operating stresses which allay irradiation embrittlement concerns.

4.07.1 Introduction

Silicon carbide (SiC) has been studied and utilized in nuclear systems for decades. Its primary use was, and still is, as the micro pressure vessel for high — temperature gas-cooled reactor fuels. For these so-called TRI-ISOtropic (TRISO) fuels, the SiC is deposited via a gas-phase decomposition process over two layers of pyrolytic graphite surrounding the fuel kernel. In addition to being strong enough to with­stand the pressure buildup from the fission product gas liberated, this SiC layer must also withstand chemical attack from metallic fission products such as palladium and the mechanical loads derived from irradiation-induced dimensional changes occurring in the pyrolytic graphite. More recent nuclear appli­cations of SiC include its use as structural composites

(i. e., SiC/SiC) for high-temperature gas-cooled reac­tors and for fusion power systems. The possibility of using composite and monolithic SiC thermal insula­tors for both fusion and fission systems is also being investigated. Moreover, both monolithic and compos­ite forms of SiC are being investigated for use in advanced sodium fast, advanced liquid salt-cooled, and advanced light water reactors.

In this chapter, the effects of neutron irradiation on relatively pure, radiation resistant forms of SiC are discussed. This chapter has been limited to the effects of irradiation on the microstructure, and the mechanical and thermal properties of SiC, although it is recognized that environment aspects such as oxida­tion and corrosion will also be factors in eventual nuclear application. These areas are not discussed here.

Cracks that occur between the liquidus and the soli­dus are commonly termed ‘hot cracks.’ Hot cracks can be further differentiated depending on whether they occur in the composite region of the weld bead on cooling (solidification-type) or in the partially melted or heat-affected zone where a composition gradient or low melting phase acts to locally depress the solidus of the alloy (liquation-type). Additionally, a third type of‘hot cracking,’ that is, hot tearing, can be
distinguished from solidification and liquation cracks. Hot tears are primarily mechanical in nature, driven by the geometry and stresses in the weldment. A schematic of the locations and representative micrographs of the different types of supersolidus cracking are shown in Figure 2.

In a nuclear reactor, high energy, fast neutron flux leads to the displacement of carbon atoms in the graphite crystallites via a ‘cascade.’ Many of these atoms will find vacant positions, while others will form small interstitial clusters that may diffuse to form larger clusters (loops in the case of graphite) depending upon the irradiation temperature. Con­versely, vacancy loops will be formed causing the lattice structure to collapse. These vacancy loops will only become mobile at relatively high tempera­tures. The production of transmutation gas from impurities is not an issue for highly pure nuclear graphite, as the quantities of gas involved will be negligible and the graphitic structure is porous.

The change in graphite properties is a function of the displacement of carbon atoms. The nature and amount of damage to graphite depends on the partic­ular reactor flux spectrum, which is dependent on the reactor design and position, as illustrated in Figure 6.

It is impractical to relate a spectrum of neutron energies to a dimensional or property change at a

single point in a material such as graphite. Therefore, an ‘integrated flux’ is used and is discussed later.

4.11.5.1 Early Activation Measurements on Foils

Although one cannot directly measure the damage to graphite itself, it is possible to measure the activation of another material, because of nuclear impacts adja­cent to the position of interest. This activation may then be related to changes in graphite properties.

This was done in early experiments using cobalt foils and by measuring the activation arising from the 59Co(n, g)60Co reaction. This reaction has a cross-section of 38 barns and 60Co has a half-life of 5.72 years, which need to be accounted for in the fluence calculations. Such foils were included in graphite experiments in BEPO and the Windscale Piles, and are still used today for irradiation rig validation and calibration purposes.

In these early experiments, after removal from the reactor, cobalt foils were dissolved in acid, diluted, and the decay rate measured. A measure of fluence could then be calculated from knowledge of the following:

• the solution concentration

• the time in the reactor

• the decay rate

• the activation cross-section

Source: Simmons, J. H. W. The Effects of Irradiation on Graphite; AERE R R 1954; Atomic Energy Research Establishment, 1956.

Unfortunately the 59Co(n, g)60Co reaction is mainly a measure of thermal flux and atomic displa­cements in graphite are due to fast neutrons. An improvement was the use of cobalt/cadmium foils, but this was not really satisfactory. Measurements made in this way are often given the unit, neutron velocity time (nvt).

Table 3 gives an example of thermal flux deter­mined from cobalt foils defined at a standard posi­tion in the center of a lattice cell in BEPO. Graphite damage at other positions in other reac­tors could then be related to the standard position in BEPO.

Irradiated creep experiments were carried out between 350 and 650 °C on both PGA and Gilsocarbon graph­ite.91 Some low fluence experiments were also carried out in Calder Hall86 which were used to define the so-called ‘primary creep.’ Creep strain data ecr was normalized to elastic strain units (esu) by dividing by the applied stress (s) and multiplying by the

unirradiated SYM (E0) as given below:

E0 ecr

esu =

а

Surprisingly, they found that by doing this, the creep data for these two types of graphite, with very differ­ent microstructures, could be fitted to a simple ‘creep equation’ of the form

аа

ecr = [exp(-4g)] + °.23 g [54]

E0 E0

where g is the fast neutron dose. This is illustrated in Figure 57.

The primary creep strain is assumed to be recov­erable on removal of the load while still under irradi­ation. Some evidence for this came from out-of-pile measurement experiments such as the FLACH experiments.9 However, if the specimens had been left unloaded for longer duration, more than 1 esu may have been recovered. In addition to this, an experiment carried out on precrept samples, that is,

samples of PGA and Gilsocarbon irradiated in a creep experiment in the BR-2 reactor and then irra­diated with the load removed in DIDO and DFR respectively, exhibited more than 1 esu (a recovery in the region of 6 esu in the case of Gilsocarbon in DFR).

4.01.1.4.1 Crystalline to amorphous transformation of Zr-(Fe, Cr, Ni) intermetallic precipitates

In addition to point-defect cluster formation, irra­diation of metals can affect the precipitation state as well as the solid solution. In the case of zirco­nium alloys, while investigating the effect of irradi­ation on corrosion, TEM observations revealed that for Zircaloy, irradiated at temperatures typical for commercial light water reactors (lower than 600 K), Zr(Fe, Cr)2 precipitates began to become amorphous after a fluence of about 3 x 1025 nm~2. Interestingly, the other common precipitate in Zy-2, Zr2(Fe, Ni), remained crystalline up to higher irra­diation doses.77 The instability of these precipitates under irradiation is of great importance since the secondary-phase precipitate plays a major role on

(b)

(d)

Figure 8 Examples of radiation-induced cavities in zirconium alloys. (a) Annealed crystal-bar zirconium, prism foil, 673K, 1.2x1025n/m2; (b) annealed zircaloy-2, prism foil, 673K, 1.2x1025n/m2; (c) annealed Zr-2.5 wt% Nb, basal foil, 923K, 0.7x1025n/m2; (d) typical cavity attached to inclusion on a grain boundary, material (c). Adapted from Gilbert, R. W.; Farrell, K.; Coleman, C. E. J. Nucl. Mater. 1979, 84(1-2), 137-148.

the corrosion resistance of Zircaloy (see Chapter 5.03, Corrosion of Zirconium Alloys).

The effect of temperature on the crystalline to amorphous transformation has been studied by vari­ous authors.7 ’ 3 It is shown that at low tempera­

tures (353 K), under neutron irradiation, both Zr(Fe, Cr)2 and Zr2(Fe, Ni) undergo a rapid and complete crystalline to amorphous transformation. As the irra­diation temperature increases, a higher dose is required for amorphization. It is indeed seen that, at 570 K, Zr(Fe, Cr)2 precipitates undergo only a partial amorphous transformation and Zr2(Fe, Ni) particles remain crystalline (Figure 9).

It is also observed that the crystalline to amor­phous transformation starts at the periphery of par­ticles, and then the amorphous rim moves inward until the whole precipitate becomes fully amor­phous. The chemical concentration profile within the precipitates also exhibits two distinct zones corresponding to the two different states: the crystal­line core and the amorphous periphery. It is observed that the amorphous layer exhibits a much lower iron

500
400
FeK,
300
CrK.
O. О
200
100
0 350
450 550 650 750 850 Energy (eV)x10-1
500
CrK,
‘a
0.1 mm

350 450 550 650 750 850

Energy (eV)x10-1

Figure 10 Crystalline to amorphous transformations of Zr(Cr, Fe)2 particle in Zy-4 irradiated at 560 K at 3.5 x 1025 n m~2. EDX spectrum shows that the amorphous volume is coincident with a depletion of Fe. Adapted from Griffiths, M.;

Gilbert, R. W.; Carpenter, G. J. C. J. Nucl. Mater. 1987, 150(1), 53-66.

content than the precipitate, the iron profile showing a local drop from the standard value of 45 at.% to below 10 at.% (Figure 10).

At higher temperatures (T> 640 K), amorphiza — tion was not detected and the precipitates remain crystalline, but some authors79 have nevertheless observed loss of iron and even total dissolution of Zr2(Fe, Ni) and Zr(Fe, Cr)2 precipitates and redistri­bution of alloying elements.

The crystalline to amorphous transformation is eas­ily understood in terms of ballistic radiation-induced disordering at a temperature where recombination ofpoint defects or recrystallization within the interme­tallic precipitate is too slow to compensate for the rate of atomic displacement (at 350 K).79 The dissolution of alloying elements remains limited at this low tempera­ture and the amorphization is mainly due to sputtering, that is, transfer of material from the particle because of atomic displacements by neutrons. When the point — defect concentration becomes too high and/or when the chemical disordering is too high, the crystalline structure is destabilized and undergoes a transforma­tion to an amorphous phase.75,79

The fact that the Zr2(Fe, Ni) phase remains crys­talline at intermediate temperatures (520-600 K) is presumably due to a more rapid reordering than the disordering in this structure (Zintl phase structure).

Concerning the Zr(Fe, Cr)2 (Laves phase structure), it is seen that the amorphization starts at the precipi­tate-matrix interface forming a front that gradually moves into the precipitate. The amorphization is believed to happen by a deviation from stoichiometry due to a ballistic interchange of iron and zirconium atoms across the precipitate-matrix interface. It also agrees with the observed kinetics of amorphization, predicting an amorphous thickness proportional to fluence and the absence of an incubation period for the transformation to start.84

The reason for the depletion of iron from the precipitates is not clearly understood yet, according to Griffiths et al.79 It is suggested that iron may be in some form of irradiation-induced interstitial state in irradiated Zr-alloys and may then diffuse intersti­tially out of the intermetallic particles.

At high temperatures (640-710 K), corresponding to 0.3 Tm the thermal activation is sufficient to induce dynamic recrystallization impeding the amorphiza — tion of the precipitates. However, depletion and some precipitate dissolution would still occur, but the level of damage necessary for amorphization would not be reached due to the absence of cascade damage.8 Because of the high mobility of Fe and Cr, redistribution of solute can occur, leading to secondary-precipitate formation.

While irradiation creep is relatively well understood the effect of radiation on thermal creep and thereby

creep rupture is not as well defined. In general it appears that creep rupture properties are not improved by irradiation and are adversely affected as shown in the example of Figure 87.1 ,

As shown in Figure 88 Ukai and coworkers have compared the reduction in rupture life in air, sodium, and after irradiation in FFTF, demonstrating that the largest influence is due to irradiation.193 There is some evidence that irradiation in neutron spectra that pro­duce high He/dpa ratios will decrease rupture life, especially at higher temperatures, compared to irradi­ation in fast reactors due to the accumulation ofhelium bubbles on grain boundaries and triple points.191,192

It is possible to improve the in-reactor stress rupture properties of a given steel by additions of selected trace elements such as P and B, both of which are known to affect the distribution and stability of carbide phases. An example is shown in Figure 89.194 Fortuitously, such additions also add to the swelling resistance of such steels.

Experimental information on the nature of matrix defects necessarily splits into evidence for clusters formed by vacancies and interstitial point defects. Significant insight into potential vacancy clusters has come from positron studies of irradiated and model steels. Positrons are well established as a tool for probing vacancy-type defects in solids, where the defect concentration is typically 1018cm~3. (Posi­trons are attracted to regions of the lattice which are more ‘open’ than average. The most obvious such regions are vacancies and vacancy clusters (the larger clusters being stronger positron traps). Less obvious open regions include the tensile parts of a dislocation strain field (even around an interstitial loop) or incoherent particle-matrix interfaces. The positron annihilation techniques have included both positron lifetime (PA-t) and lineshape (PALA) ana­lyses or, more recently, CDB. The latter is an impor­tant recent development. CDB (also known as positron annihilation spectroscopy-orbital angular momentum distribution, PAS-OEMS) measures the energy spectrum of the gamma rays produced at the positron annihilation sites, and thus the momenta of the electrons at those sites. The energy (or momentum) spectra characterize the elements sur­rounding the positron traps.72

Overall, the experimental data provide strong evi­dence for open-volume clusters that are sensitive to positrons, and for PA being a useful technique for studying MD (see, e. g., Buswell and Highton,73 Dai et a/.,74 Carter et a/.54). In model alloys, a number of authors report evidence for large vacancy clusters or microvoids after irradiation. Most interestingly, although the positron lifetime studies have indicated the presence of vacancy clusters in model alloys, such studies have not identified vacancy clusters in com­mercial steels (e. g., Dai et a/.74). The study by Valo eta/.75 is most convincing in this aspect, as the authors investigated both model alloys and RPV steels.

Evidence for interstitial clusters has come primar­ily from studies of model alloys or steels irradiated to very high doses in excess of that of interest for most power reactor applications. Such studies have nor­mally employed TEM techniques that are sensitive to the strain field associated with small dislocation loops. The imaging of such features is difficult in ferritic materials because of the need to correct for the image distortion caused by the ferromagnetic behavior of the samples, and because of the contrast arising from surface oxide on the thinned specimens. Critically, the resolution for small dislocation loops is ^2-3 nm in even well-prepared samples imaged in higher voltage microscopes. Krishnamoorthy and Ebrahimi76 and Hoelzer and Ebrahimi77 reported the formation of visible interstitial loops in Fe irra­diated to 4.63 x 1023 nm~2; E> 1MeV at ^280 °C; the loops increased in size and decreased in number density after annealing at 500 °C.

There have been fewer studies on irradiated steels, but in MnNiMo steels little evidence for dislocation loops has been reported. Soneda and coworkers have undertaken weak-beam TEM obser­vations of RPV surveillance steels containing 0.06 and 0.12 wt% Cu irradiated to 4x 1019 ncm~2.65

Soneda reported the formation of interstitial dislo­cation loops, whose diameter and number density are 2-3 nm and of the order of 1021m — , respec­tively. At very high doses (and dose rates) a uniform density of loops has been observed. Fuji and Fukura78 undertook a weak-beam TEM study for MD in A533B RPV steel produced by 3 MeV Ni2+ ion irradiation to a dose of 1 dpa at 290 °C. The MD was found to consist of small dislocation loops. The observed and analyzed dislocation loops have Burgers vectors b = a <100> (Figure 13). The dis­location loops have a mean image size d = 2.5 nm and the number density is about 1 x 1022m-3. Most of the loops are stable after thermal annealing at 400 °C for 30 min. This indirect evidence suggests that their nature is interstitial.

Kocik et al..79 examined the radiation damage microstructures in Cr-Mo-V surveillance base metal and weld containing (^0.06-0.07 wt% Cu and 0.012-0.014 wt% P) irradiated up to 6 x 1024nm-2 (E> 1 MeV) for times up to 5 years at 265 °C. TEM examination of the irradiated materials revealed, in both the base metal and the weld metal, black dots, small (resolvable) dislocation loops, and small precipitated particles. Clouds of defects are formed along dislocations at higher neutron fluences, and it was only at the higher fluence that loops that may not be associated with dislocations could be seen. Interactions were observed between defects and (as-grown) dislocations that result in a rebuild of dislocation substructure. Miller et a/.80 examined the radiation damage microstructures in similar Cr-Mo-V surveillance base metal and weld. They reported manganese-, silicon-, copper-, phosphorus-, and carbon-decorated dislocations and other features in the matrix of the neutron-irradiated base and weld materials.

4.05.4.5.2 Development with flux and fluence and irradiation temperature

The most important inference from the mechanical test data is that hardening and embrittlement are proportional to the square root of fluence in low copper steels. Early theoretical and experimental work by Makin and coworkers81,82 demonstrated that a square root dependency on dose was consistent with the hardening arising from the cutting, by glide dislocations, of irradiation-produced obstacles, and that in the early stages of irradiation the number density of clusters is proportional to the irradiation exposure. Thus, in irradiated low Cu RPV steels, there is continuous production of hardening centers during irradiation. Further, the linear dependence of hardening on irradiation temperature from 150 to ^300 °C in CMn steels and low Ni A533B weld­ments implies that thermal stability of MD clusters is

important.83

There are relatively few studies that generate insight into the effect of flux and fluence on MD itself. Unsurprisingly, studies of model alloys tend to emphasize the increase in number density (and size) of the vacancy-rich clusters with increasing dose. Kampmann et a/.84 found void-like features 1-2 nm diameter in Cu-free ternary Fe-Ni-P/Mn alloys irradiated 2-25 x 1018ncm-2. The authors considered that the microvoid numbers increase with dose up to ^5 x 1018ncm-2, and then either remain constant or decrease. Analyzing positron annihilation data from annealing studies of neutron — irradiated A533B plate, A508-3 forging, and welds, Carter et al.54 considered that increasing the dose from 1 x 1018to20 x 1018ncm-2 at 290 °C increased the volume fraction of vacancy clusters, probably via increasing both the number density and average size of the clusters. Increasing the flux from 6 x 1011 to 5 x 1012 n cm-2s-1 increased either the number den­sity or the mean radius, probably the radius.

Postirradiation annealing has been shown to be a powerful means ofinvestigating the nature ofthe MD further. A major development has been characterizing the matrix defect term as being due to two compo­nents; first, stable matrix defects (SMD) and second, at high fluxes, unstable matrix defects (UMD) (see, e. g., Mader et al.85). UMD are matrix defects that, although thermally unstable at the irradiation temperature, are frozen into the microstructure during the cooldown after irradiation. Such studies have also established that MD and hardening of low Cu steels will be dose rate dependent at high dose rates (>1-5 x 1012 n cm-2 s-1, E > 1 MeV).85

Soneda65 modeled the effects of dose, dose rate, and irradiation temperature on the defect accumula­tion in bcc-Fe using the kinetic Monte Carlo (KMC) method.65,86 Jones and Williams83 proposed a model that describes the irradiation temperature depen­dence of the embrittlement of low Cu materials, AT = a x FT x (‘t)1/2, where AT, a, and ‘t are the transition temperature shift (TTS), constant coefficient, and dose, respectively, and FT = 1.869­4.57 x 10~3T (°C). This model was studied using a KMC simulation. The number densities of both vacancies and self-interstitial atom (SIA) clusters exhibited a linear temperature dependence with a slope equivalent to that of FT, and Soneda considered that the origin of the form of the FT term can be understood from the temperature dependence of point defect cluster formation.

4.08.3.2.1 Mechanically alloyed powder characterization

The computed phase diagram of the Fe-0.13C-2W — 0.2Ti system without Y2O3 is shown in Figure 6 with respect to carbon content. For a carbon content of 0.13 wt%, a single austenite g-phase containing TiC carbide exists at a normalizing temperature of 1050 °C. The equilibrium g/g + 8-phase boundary at this temperature corresponds to a carbon content of 0.08 wt%, beyond which 8-ferrite is not stable. The specimens without and with 0.1 wt% Y2O3 exhibit the full martensite structure, whereas the specimens with 0.35 and 0.7 wt% Y2O3 exhibit a dual phase compris­ing both martensite and ferrite phases. Digital image analyses show that the area fraction of the ferrite phase is ^0.2 for specimens with 0.35 and 0.7 wt% Y2O3. High-temperature X-ray diffraction measurement at 950 °C showed a considerable difference; the specimen without Y2O3 shows diffraction peaks that correspond only to the austenite g-phase, whereas specimens with 0.35 and 0.7 wt% Y2O3 show diffraction peaks corresponding not only to an austenite g-phase but to a
ferrite phase as well. The austenite g-phase transforms to the martensite phase, but the ferrite phase remains unchanged by quenching. Considering that the ferrite phase is formed only in the specimens containing 0.35 and 0.7 wt% Y2O3, and that four types of ODS steels have an identical chemical composition except for Y2O3 content, the Y2O3 particles could suppress the a—g reverse transformation.

Figure 722 shows the results of dilatometric mea­surement when 9Cr-0.13C-2W-0.2Ti is heated without and with 0.35 wt% Y2O3. In the case of the specimen without Y2O3, the linear thermal expansion begins to decrease from an Ac1 point of850 °C to an Ac3 point of 880 °C, due to the reverse transformation of a-g-phase, which corresponds reasonably well with the computed phase diagram. The addition of 0.35 wt% Y2O3 induces an increase up to an Ac3 point of 935 °C. By comparing both curves, it was found that the specimen with 0.35 wt% Y2O3 exhibits a smaller degree of reduction in linear thermal expansion during the reverse transformation of the a-g-phase; this obser­vation indicates that the entire a-phase could not be transformed to a g-phase. This untransformed ferrite phase was designated as a residual ferrite.

4.09.4.1 Low-Alloy Steels

Low-alloy steels generally have good weldability with the main concern being liquation cracking near impurities (Figure 8), the propensity of some grades (esp. Cr-Mo steels and some pressure vessel grades) to reheat-type cracking (discussed earlier) and to hydrogen-induced cracking.1,91,92 The suscep­tibility to hydrogen-induced cracking is controlled by four major considerations:

1. The composition of the steel

2. The mobile hydrogen concentration

3. The stresses in the weldment

4. The thermal management of the weld

In general, the more hardenable the steel (i. e., the more easily martensite is formed), the more suscep­tible it is to cracking. Since hardenability generally increases with carbon content and alloying additions, several parameters have been developed to gauge susceptibility to hydrogen-induced cracking, includ­ing the carbon equivalent (Ceq) in eqn [2] and the Ito- Bessyo ‘cold cracking’ parameter (Pcm) (eqn [3]).93 The concentrations in eqns [2] and [3] are for weight percentage.

Cr + Mo + V Mn Ni + Cu

——— 5——— + — + — І5- и

V Mo

5 x B + C + To+15 +

Si Ni

+ 30 + 60

The hydrogen concentration in the weldment can be minimized by proper cleanliness, preheat, and post­weld heat treatment.92,94 Additionally, microstruc­tural hydrogen traps can provide significant benefit by preventing hydrogen redistribution to regions of high stress.37 As the mobile hydrogen concentration is most detrimental, significant work has gone into standards to accurately assess the amount of diffusible hydrogen in steels.37,92,94 Minimizing residual stres­ses and avoiding geometric stress concentrators (e. g., notches) in the weldment also impart resistance to hydrogen-induced cracking.

Thermal management of the weldment (e. g., pre­heating, interpass temperature, bead tempering, and postweld heat treatment) is also critical to the mitiga­tion of hydrogen-induced cracking. Preheating of components and interpass temperature control act to outgas hydrogen or hydrogen-bearing compounds (e. g., water) and lower the cooling rate. Additionally, careful control of heat input can produce hydrogen — resistant microstructures (i. e., bead tempering). Postweld heat treatment can act to lower residual stresses, produce beneficial hydrogen traps, and remove dissolved hydrogen from the weld.3,4,95,96 Several sources provide guidelines to mitigate hydro­gen cracking in specific grades of steel.95-97