In reviewing radiation damage effects in ferritic steels, it is important to recognize that a range of ferritic steels have been employed in commercial reactors. Such steels were necessarily employed in thick sec­tions and the fabrication of the vessels involved weld­ing of preformed plates or forgings (for a description of typical vessels, see Steele and Sterne1). Frequently, in early designs, the welds were located opposite the center of the reactor core and received the highest neutron dose. Commercially available ferritic steels were employed in the construction ofthe first reactors (both for Magnox and LWR designs).

For LWRs, which have to contain higher pressure than the pressure vessels of Magnox reactors, the desirability of using steels of high toughness, ade­quate strength, and weldability in thick sections, combined with good service experience, has nar­rowed the choice to a number of low alloy steels; for example, those containing manganese, nickel, and
molybdenum. Indeed, the ASME specification for quenched and tempered vacuum-treated carbon and alloy steel forgings for pressure vessels permits only SA-508 Class 2 and Class 3 compositions. Early ves­sels were constructed from A302 and A302B plates. Vessels constructed in Russia (so-called VVER reac­tors) are fabricated from high-strength CrMoV steels.

The steels used in vessel construction of Magnox pressure vessels generally consisted of simple C-Mn ferritic plates, either Si-killed or Al-grain refined, in the normalized and stress-relieved condition. Table 12 shows the average, or range of, chemical compositions for the plates, manual metal-arc (MMA) welds, submerged-arc (SMA) welds, and forgings, respec­tively, as derived from the populations of vessel cut-outs ex-construction, surveillance samples, and contemporary reproduction vessel materials.

Compositions of commercial MnMoNi steels used for modern LWR pressure vessels are given in Table 2 9-12 A detailed description of the develop­ment of modern pressure vessel steels can be found in

Druce and Edwards.13 It is to be noted that commer­cial steels are complex with many elements present. Although the concentrations of additional elements are low, all have the potential to influence the radia­tion damage response of the steel.

The deleterious effects of residual impurities such as Cu and P on the in-service degradation in properties (see Section 4.05.3) have been increas­ingly recognized. Consequently, over time, the levels of Cu and P in ferritic pressure vessel steels have been reduced and in modern steels, such as the for­gings and welds employed in fabricating the vessel for the PWR at Sizewell in the United Kingdom, the levels of such impurities are well-controlled (see Table 3).

Finally, it is important to recognize that these steels have a complex microstructure at start of life (SOL), and that there may be significant spatial vari­ation in the local microstructure.14 For example, the local dislocation density or size and number of second-phase precipitates may vary significantly both within a given material and also between plates, welds, or forgings produced to the same nominal specification.

According to Lee et a/.,43 the effect of neutron irradi­ation on the specific heat of SiC was negligibly small. The specific heat of SiC is therefore assumed to be unchanged by neutron irradiation, although this has not been verified at high dose. A single study5 also indicated that stored energy (Wigner energy) occurs in SiC irradiated in the point defect regime. The relative amount of stored energy appears to be less than that of graphite.44

Because of a low density of valence band electrons, thermal conductivity of most ceramic materials, SiC in particular, is based on phonon transport. For a ceramic at the relatively high temperature associated with nuclear applications, the conduction heat can be generally described by the strength of the individual contributors to phonon scattering: grain boundary scattering (1/Kgb), phonon-phonon interaction (or Umklapp scattering 1/Ki), and defect scattering (1/Kd). Because scattering of each of these types occurs at differing phonon frequencies and can be considered separable, the summed thermal resis­tance for a material can be simply described as the summation of the individual components; that is, 1/K = 1/Kgb + 1/Ku + 1/Kd. As seen in Figure 7, the unirradiated thermal conductivity of SiC is highly dependent on the nature of the material (grain size, impurities, etc.) and the temperature. While materials can be optimized for low intrinsic defect, impurity,

and grain boundary scattering, the temperature — dependent phonon scattering cannot be altered and tends to dominate at high temperature (above about 673 K for SiC).

The effect of irradiation on SiC in the tempera­ture range of ^423-1073 K (the point defect regime) is to produce simple defects and defect clusters that very effectively scatter phonons. For ceramics pos­sessing high thermal conductivity, the irradiation — induced defect scattering quickly dominates, with saturation thermal conductivity typically achieved by a few dpa. Moreover, as the irradiation-induced defect scattering exceeds the phonon-phonon scat­tering, the temperature dependence of thermal con­ductivity is much reduced or effectively eliminated.

The rapid decrease as well as saturation in thermal conductivity of CVD SiC upon irradiation in the point-defect regime has been reported by several authors.8,34,45,52,53 Figure 8 shows this rapid decrease in thermal conductivity for fully dense CVD SiC, including new data, previous data from the authors,52,53 and that of Rohde.45 It is noted that the data of Thorne is omitted as the material was of exceptionally low density for a CVD SiC material. Moreover, the data of Price34 is published with a range of fluence that is not valuable in the figure.

In recent papers by Snead on the effects of neu­tron irradiation on the thermal conductivity of cera — mics,53 and specifically on SiC,16,52 the degradation in thermal conductivity has been analyzed in terms of the added thermal resistance caused by the neutron irradiation. The thermal defect resistance is defined as the difference between the reciprocals of the irradiated and nonirradiated thermal conductivity (1/Kd = 1/Kirr-1/Knonirr). This term can be related directly to the defect type and concentration present in irradiated ceramics.53 Moreover, this term can be used as a tool to compare the thermal conductivity degradation under irradiation of various ceramics or, for example, various forms of SiC. It has been shown that, for certain high purity forms of alumina, the accumulation of thermal defect resistance is very similar even though the starting thermal con­ductivities of the materials are substantially dif­ferent. Similarly, CVD SiC was shown to have a similar accumulation of thermal defect resistance as a hot-pressed form of SiC with substantially lower (^90 W m-1 K~ ) unirradiated thermal conductivity. The utility of this finding is that if the thermal defect resistance is mapped as a function of irradiation tem­perature and dose for a form of high-purity CVD SiC, it can be applied to determine the thermal

conductivity of any high-purity CVD SiC, indepen­dent of the starting thermal conductivity. The accu­mulation in thermal defect resistance generated from the data of Figure 8 is shown in Figure 9.

Another result of the previously reported analysis on irradiated CVD SiC16,52 is that the thermal defect resistance appears to be directly proportional to the irradiation-induced swelling, although the data-set for making the previous assertion was somewhat lim­ited. A compilation plot including the previous data­set as well as the new data of Figure 9 is shown in Figure 10. It is clear from this plot that a linear relationship exists between swelling and thermal defect resistance. Moreover, there does not appear to be any effect of irradiation temperature on this result. The fact that the thermal defect resistance is proportional to the irradiation-induced swelling allows a rough estimate of thermal conductivity. As measurement of thermal conductivity for the SiC TRISO shell is not practical, while measurement of density is routine, this finding allows an indirect determination of thermal conductivity by measure­ment of the density change in the TRISO SiC shell by means of a density gradient column or some other technique.

The thermal conductivity degradation discussed up to this point has been for irradiation temperature associated with the point defect regime. For irra­diation above this temperature (the nonsaturable void swelling regime), the thermal properties are not expected to saturate (at least at low dpa). The primary reason for this is that the formation of voids and other complex defects in the high-temperature regime (which contributes to the unsaturated swelling as seen in Figure 6) contributes to phonon scattering, and these defects will not saturate. Moreover, it has been shown that the linear relationship that existed between swelling and thermal defect resistance (as seen in Figure 10) does not exist in this elevated temperature irradiation regime.16,52 This underlines the fact that the phonon scattering and swelling are not controlled by the same defects in the lower temperature ‘saturable,’ and elevated temperature ‘nonsaturable’ irradiation regimes. A compilation plot of room-temperature thermal conductivity as a function of irradiation temperature for the saturable and nonsaturable temperature regimes is given in Figure 11 .

By comparison to the unirradiated room — temperature conductivity value of ^280 W m-1 K~ ,

Fast neutron dose (x 1025nm-2 E >0.1 MeV)

it is clear that the thermal conductivity degradation in the highest temperature regimes is less dramatic, even though the swelling is rapidly increasing (see Figure 6). This is opposite to the behavior in the lower temperature, saturable regime, where high swelling corresponds to extreme reduction in thermal conductivity. Unfortunately, the data on thermal con­ductivity reduction in the nonsaturable regime is limited, and given the lack of knowledge of the spe­cific defects governing the phonon scattering, it is not possible to accurately predict behavior outside of the data-set of Figure 6.

Data presented thus far has been limited to mea­surement of thermal conductivity at room tempera­ture. As described in Figure 7, there is a dramatic dependence of thermal conductivity on measure­ment temperature. The temperature dependence of irradiated materials can be found by applying the temperature dependence of unirradiated SiC (the Umklapp thermal resistance term) to the as-neutron-degraded room-temperature values. This approximation (dashed lines) is compared to actual
data (solid lines) in Figure 12 and shows fair corre­spondence. However, it is clear that such a treatment systematically underestimated the thermal conduc­tivity degradation. This implies temperature depen­dence on the defect scattering that is not presently understood.

In contrast to solidification cracks, liquation cracks occur in the partially melted zone and the heat — affected zone and can be either interdendritic or intergranular in nature. An example of an interden­dritic liquation crack in a nickel-chromium alloy is given in Figure 2 (top), while intergranular liquation cracks in a pressure vessel steel are shown in Figure 8. The liquation cracks in Figure 8 are caused by the presence of sulfur-rich inclusions that liquate in the partially melted and heat-affected zones of the weld.

Another variation of liquation-type cracking can occur via the partial dissolution of second-phase particles, that is, the constitutional liquation mecha­nism proposed by Savage and shown experimentally by Pepe and Savage.1,16,17 In this type of cracking, the
heat from welding partially solutionizes the second — phase particles in the heat-affected zone. The result­ing concentration gradient around the particle lowers the solidus (e. g., the effect of niobium on nickel-based alloys from NbC or Ni2Nb) locally.

However, the use of R(0) is valid only for very low fluence and it should no longer be used, although one may come across its use in historic papers.

4.11.5.2 Equivalent Nickel Flux

Nickel foils were used to give a measure of the damage to graphite through the 58Ni(n, p)58Co reac­tion. This reaction has a mean cross-section of 0.107 barns and 58Co has a half-life of 71.5 days.

The change in graphite thermal resistivity was measured in the TE10 experimental hole in BEPO and the nickel flux was also measured at the same position. It was assumed that the graphite displace­ment rate f d was equal to the nickel flux f Ni at this position. For comparison, the change in graphite thermal resistivity was then measured at various other positions in BEPO, as given in Table 4

Later, the same exercise was repeated in PLUTO, the sister reactor to DIDO at Harwell, and the ratio compared to that at other positions. In this case, the ratio appeared to be largely invariant to position. Table 5 gives a few examples of the many measure­ments made.16

It was decided that the activity produced in nickel fNi could be related to the graphite damage rate by a

Modified from Bell, J.; Bridge, H.; Cottrell, A.; Greenough, G.; Reynolds, W.; Simmons, J. Philos. Trans. R. Soc. Lond. A Math. Phys. Sci. 1962, 254(1043), 361-395. b is a proportionality factor.

factor. However, care was still required with respect to the choice of reactor and irradiation location. Thus, a definition of damage based on a standard position in DIDO and a calculation route for equiva­lent DIDO nickel flux (EDNF) were devised.

It should be noted that there are difficulties related to a standard based on measurements made with nickel foils and the 58Ni(n, p,)58Co reaction because of the short half-life of 58Co and the interfering effect of the 58Co(n, g)59Co reaction. A method by Bell eta/.15 which went back to measuring activation of cobalt foils and the 59Co(n, g)60Co reaction, and then calcu­lating the ratio fNi/f Co, was used for a short while. This method used the following relationships:

115g fuel elements fNi/fco = 0.378 — 0.504b

150g fuel elements fNi/fco = 0.502 — 0.530b

where ‘b’ is the fuel burnup. However, this was not very satisfactory and it was clear that a validated calculation route was desirable, and is now becoming practicable through development in computer technology.

It is unfortunate that a validated set of graphite irra­diation creep data covering the range oftemperatures and fluences of interest for power producing reactors, as well as radiolytic oxidation in the case of carbon dioxide-cooled reactors, does not exist. In addition, there are no microstructural studies available to give an insight into the mechanism involved in irradiation creep in graphite. This has lead to much speculation and several model proposals.

Figure 59 Synergy between changes in coefficient of thermal expansion in irradiation creep specimens and change in unirradiated, stressed graphite. (a) additional change in CTE as a function of creep strain and (b) change in CTE in unirradiated stressed samples.

2.0

uj. (Л =3 "O О E 0) O)

1.6 1.2

A Й

0*_A,

0 8 — SM1-24 (axial), irradiation temperature = 850-9200C

0.0MPa (Capsule 76M-18A) О 0.0MPa (Capsule 77M-10A) 04 _ ■ 3.3MPa (Capsule 76M-18A) □ 3.3MPa (Capsule 77M-10A)

4.5MPa (Capsule 76M-18A) Д 4.5MPa (Capsule 77M-10A) 6.5MPa (Capsule 76M-18A) О 6.5MPa (Capsule 77M-10A)

0.0

0 2 4 6 8 10 12 14

Fluence (1024 ncm-2, E > 29fJ)

Figure 60 Changes in Young’s modulus in tensile crept and uncrept specimens. Reproduced from Oku, T.; Fujisaki, K.; Eto, M. J. Nucl. Mater. 1988, 152(2-3), 225-234.

4.01.2.1 Mechanical Behavior During Tensile Testing

4.01.2.1.1 Irradiation hardening: Macroscopic behavior

As for many other metals, zirconium alloys exhibit strong hardening after neutron irradiation. It is indeed observed by numerous authors90-99 and reviewed21,77,100 that the yield stress (YS), as well as the ultimate tensile strength (UTS), of both recrystallization-annealed (RXA) and stress-relieved annealed (SRA) zirconium alloys is strongly increased by neutron irradiation (Figures 12 and 13). Micro­hardness tests also prove this phenomenon.101-10 The irradiation-induced hardening increases rapidly for fluences below 1 x 1024nm~2 (E > 1 MeV), at irra­diation temperatures between 320 and 360 °C, but saturates above 1 x 1024n e

change occurs from 1 x 1024 up to 1.5 x 1025nm~2 (E > 1 MeV).92 It is however to be noticed that some authors do not find a clear saturation of the irradiation-induced hardening for fluences up to 1.5 x 1025n m~2 and irradiation temperatures between 320 and 360 °C.9297 Although the YS (and UTS) of SRA Zr alloys is significantly higher than the YS of

RXA Zr alloys before irradiation, the YS of both alloys, measured after high irradiation doses, at saturation,

21,90,100

become close.

According to Higgy and Hammad,92 and reviewed by Douglass,21 as the irradiation temperature increases from temperatures below 100 °C up to temperatures between 320 and 360 °C, the irradiation-induced hard­ening decreases. According to these authors, this shows that the accumulation of damage decreases as the irra­diation temperature increases, presumably due to recovery during irradiation.

The chemical composition seems to play a second­ary role in the irradiation-induced hardening com­pared to the effect of the metallurgical state (SRA vs. RXA). The oxygen content is nevertheless shown to have a slight effect on the irradiation-induced harden­ing. Indeed, Adamson and Bell101 have shown using microhardness tests that the irradiation-induced hard­ening is higher for RXA Zy-2 alloy with high oxygen content (1800 ppm) than in the case of an RXA Zy-2 alloy with low oxygen content (180 ppm).

It can also be noticed that the test temperature seems to have only a small influence on the irradiation — induced hardening, for a given irradiation temperature, up to a test temperature of 400 °C. Indeed, as reported by Onchi eta/.96 (Figure 14), the YS of both irradiated and unirradiated RXA Zy-2 decreases with the test temperature, the decrease being only slightly lower for the irradiated specimens between 20 and 300 °C. However, beyond a test temperature of 400 °C, a strong decrease of the irradiation hardening occurs due to the recovery of the irradiation damage.

4.03.1 Introduction

The widespread acceptance of nuclear energy depends1 on the improved economics, better safety, sustainability, proliferation resistance, and waste man­agement. Innovative technological solutions are being arrived at, in order to achieve the above goals. The anticipated sustainability, rapid growth rate, and eco­nomic viability can be ensured by the judicious choice of fast reactor technology with a closed fuel cycle option. The fast reactor technology has attained (http://www. world-nuclear. org/info/inf98.html) a high level of maturity in the last three decades, with 390 years of successful operation. The emerging inter­national collaborative projects (http://www. iaea. org/ INPRO/; http://www. gen4.org/) have, therefore, cho­sen fast reactors as one of the important constituents of the nuclear energy in the twenty-first century.

The nuclear community has been constantly striv­ing for improving the economic prospects of the technology. The short-term strategies include the development of radiation-resistant materials and extension of the lifetime of the components. The achievement of materials scientists in this field is remarkable. Three generations of materials have been developed,2 increasing the burn-up of the fuel from 45 dpa for 316 austenitic stainless steel to above 180 dpa for ferritic steels. Presently, efforts are in progress to achieve a target burn-up of 250 dpa, using advanced ferritic steels. The attempts by nuclear technologists to enhance the thermal effi­ciency have posed the challenge of improving the high temperature capability of ferritic steels. Addi­tionally, there is an inherent disadvantage in ferritic steels, that is, their susceptibility to undergo embrit­tlement, which is more severe under irradiation. It is necessary to arrive at innovative solutions to overcome these problems in ferritic steels. In the long time horizon, advanced metallic fuels and cool­ants for fast reactors are being considered for increasing the sustainability and thermal efficiency respectively. Fusion technology, which is ushering (http://www. iter. org/proj) in a new era of opti­mism with construction of the International Ther­monuclear Experimental Reactor (ITER) in France, envisages the use of radiation-resistant advanced ferritic steels. Thus, the newly emerging scenario in nuclear energy imposes the necessity to reevalu­ate the materials technology of today for future applications.

The genesis of the development of ferritic steels is, indeed, in the thermal power industry. The develop­ment of creep-resistant, low alloy steels for boilers and steam generators has been one of the major activities in the last century. Today, the attempt to develop ultra super critical steels is at an advanced stage. Extensive research of the last century is responsible for identifying certain guidelines to address the concerns in the ferritic steels. The merit of ferritic steels for the fast reactor industry was established3 in the 1970s and since then, extensive R&D has been carried out4 on the application of ferritic steels for nuclear core component.

A series of commercial ferritic alloys have been developed, which show excellent void swelling resis­tance. The basic understanding of the superior resistance of the ferrite lattice to void swelling, the nature of dislocations and their interaction with point defects generated during irradiation have been well understood. The strengthening and deformation mechanisms of ferrite, influence of various alloying elements, microstructural stability, and response of the ferrite lattice to irradiation temperature and stress have been extensively investigated. The mechanism of irradiation hardening, embrittlement and methods to overcome the same are studied in detail. Of the dif­ferent steels evaluated, 9-12% Cr ferritic-martensitic steels are the immediate future solution for fast reac­tor core material, with best void swelling resistance and minimum propensity for embrittlement.

The high temperature capability of the ferritic steels has been improved from 773 to 973 K, by launching the next generation ferritic steels, which are currently under evaluation for nuclear applica­tions, namely the oxide dispersion strengthened (ODS) ferritic steels (see Chapter 4.08, Oxide Dis­persion Strengthened Steels). Conceptually, this series of steels combines the merits of swelling resis­tance of the ferrite matrix and the creep resistance offered by inert, nanometer sized, yttria dispersions to enhance the high temperature limit of the ODS steels to temperature beyond 823 K. The concerns of this family of materials include optimization of the chemistry of the host lattice, cost effective fabrication procedure, and stability of the dispersions under irra­diation, which will be discussed in this article.

The present review begins with a brief introduc­tion to the basic metallurgy of ferritic steels, summar­izing the influence of chemistry on stability of phases, decomposition modes of austenite, different types of steels and structure-property correlations. The main thrust is on the development of commercial ferritic steels for core components of fast reactors, based on their chemistry and microstructure. Hence, the next part of the review introduces the operating conditions and radiation damage mechanisms of core compo­nents in fast reactors. The irradiation response of ferritic steels with respect to swelling resistance, irra­diation hardening, and irradiation creep are high­lighted. The in-depth understanding of the damage mechanisms is explained. The main concerns of fer­ritic steels such as the inferior high temperature irra­diation creep and severe embrittlement are addressed. The current attempts to overcome the problems are discussed. Finally, the development of advanced creep-resistant ferritic steels like the ODS steels, for fission and fusion applications are presented. The application of ferritic steels for steam generator cir­cuits and the main concerns in the weldments of ferritic steels are discussed briefly. The future trends in the application of ferritic steels in fast reactor technology are finally summarized.

The small defects formed in irradiated steels and model alloys can act as barriers to dislocation move­ment and therefore result in an increase in yield strength and hardness. Particularly important is the hardening from the copper-enriched precipitates/ clusters formed during irradiation in the high copper steels which can be modeled using the Russell-Brown model.94 The Russell-Brown model of hardening due to copper precipitates is a modulus interaction the­ory, based on the reduction in energy of the segment of dislocation, which passes through a relatively soft copper particle in the iron matrix. As the energy of the dislocation is proportional to the modulus of the host material, an attractive force will act on the dislo­cation because the modulus of copper is less than that of iron. Russell and Brown estimated the attractive force as a function of copper volume fraction, and demonstrated that this could adequately describe hardening in Fe-Cu alloys. A key element in applying the Russell-Brown model is the estimation of the modulus. Three approaches have been employed, using the modulus for fcc Cu, or computing values for bcc Cu,95 or fitting to experimental data. The last one is the most common approach. The matrix hard­ening may be estimated from the response of low Cu steels (Cu < 0.1 wt%).

The individual hardening contributions from CECs and the MD must be combined with one another, as well as with the hardening from the pre­existing microstructure. The limiting rules for such superposition are a linear sum (LS) law and a square root of the sum of the squares (RSS) law.96 Computer models can be employed to determine the exact superposition law to be employed.97,98 Figure 14 shows a scatter plot, where the measured Asy is compared to the predicted values.99 It can be seen that excellent agreement can be achieved.

Bacon and Osetsky100 carried out molecular static (MS) and molecular dynamics (MD) simulations of the passage of a dislocation through a bcc Cu precip­itate. The MS simulations led to a dependence of hardening on precipitate size which differed from that predicted by the Russell-Brown model. How­ever, Odette (see Section 2 of Eason et at29) found that the Russell-Brown model gave slightly better agreement with the experimental data.

It should be added that further insight into the parameters controlling the hardening is obtained from CECs by combining microstructural data with

mechanical property data (particularly hardening or yield stress increase) where the MD has been sub­tracted from the total measured increase.

Nanoindentation measurements were conducted in order to evaluate the mechanical properties of the residual ferrite itself. The trace of a Berkovich tip can be placed within the interiors of the residual ferrite regions, while conventional micro-Vickers diamond tips using 100-mN loads cover 7 x 7 pm. Figure 10 shows the hardness change in the individual phases measured by this nanoindentation technique as a

Figure 10 Hardness change at room temperature as a function of tempering conditions for the residual ferrite and tempered martensite. NT: normalizing and tempering; FC: furnace cooling. Ukai, S.; Ohtsuka, S.; Kaito, T.; Sakasegawa, H.; Chikata, N.; Hayashi, S.; Ohnuki, S. Mater. Sci. Eng. A 2009, 510-511,115-120.

parameter of the tempering conditions.26 The decrease in hardness is significantly restricted in the residual ferrite as compared to that of the martensite phase in terms of increasing the tempering conditions. The overall hardness measured by the micro-Vickers tester is also shown by the broken line which covers both the residual ferrite and martensite, therefore, representing the average hardness of both phases. Hardness Hv is correlated with yield stress sy using the relationship provided by Tabor.27 For tempering conditions at 800 °C for 58 h, which is equivalent to tempering at 700 °C for 10 000 h based on the LMP (Larson-Miller parameter), hardness can be converted to yield stress at room temperature for the individual phases: 1360 MPa for the residual ferrite and 930 MPa for the tempered martensite. The yield strength of the residual ferrite is 1.5 times higher than that of mar­tensite at tempering at 700 °C for 10 000 h.

A full ferrite ODS steel and full martensite ODS steel were manufactured, and the oxide particle dis­tribution in both ODS steels was measured by TEM. The results are shown in Figure 11.28 It is obvious that a few nanometer-sized oxide particles are finely distributed in the full ferrite ODS steel, whereas their size is coarsened in the bi-modal distribution in the martensite ODS steel. Considering that the residual ferrite phase belongs to full ferrite ODS steel, resid­ual ferrite contains fine (nanosized) oxide particles which are responsible for higher strength in residual ferrite containing ODS steels. In regard to the bi­modal distribution of oxide particles in martensite

ODS steels, the a-g-phase transformation could induce the coarsening of oxide particles by disturbing the interface coherency between these particles and the g-phase matrix.

Like austenitic stainless steels, single-phase nickel — based alloys are susceptible to solidification and liquation cracking (Chapter 2.08, Nickel Alloys: Properties and Characteristics). Common melting point suppressants that must be controlled to avoid cracking include tramp elements such as sulfur and intentional alloying additions such as boron, zirconium, niobium, and molybdenum.7,10-15,18,109,110 Addition­ally, both low-strength and precipitation hardenable grades can be susceptible to PIC mechanisms (ductility dip and strain-age cracking) as discussed previously.

4.09.4.2 Zirconium Alloys

Zirconium alloys are readily weldable but as with all reactive metals, need special precautions to prevent pickup of interstitial elements such as oxygen, carbon, and nitrogen that can degrade both the mechanical

properties and the corrosion performance of the weld (Chapter 2.07, Zirconium Alloys: Properties and Characteristics).9 , , 2 Vacuum or inert gases

(argon, helium, or Ar-He mixtures) can be used to shield zirconium, but again, care needs to be taken to ensure sufficient vacuum level or gas purity to pre-

92,111

vent contamination.

Zirconium alloys can be susceptible to both super­solidus and subsolidus (i. e., hydride-type) cracking. Supersolidus cracking is typically solidification-type and many common alloying elements and/or poten­tial contaminants promote susceptibility. For exam­ple, iron, nickel, chromium, and copper all stabilize low-temperature eutectic reactions, and small con­centrations can greatly increase the solidification temperature range. An example of this is shown in Figure 22, which compares the maximum crack length in transvarestraint tests for a Zr-Cr alloy, Zircaloy-4 (Zr-4), and Zr-2.5Nb welded under iden­tical conditions. As shown, the Zr-Cr alloy exhibits the most susceptibility to solidification cracking.

Acknowledgments

The Authors are indebted to the welders, technicians, specialists, and engineers of the Welding & Materials Process Development Unit at Knolls Atomic Power

E

E,

-C

D>

C

Ф

О

2

о

g

2

Laboratory, whose dedication and expertise made this work possible. They are also grateful to Dr. David

S. Knorr of General Electric for his important con­tributions to the manuscript.