An ambitious target of increasing the temperature and pressure of steam in many power plants has provided a high impetus for the development of steels with better high temperature properties. Very often, the weld joints play a crucial life limiting role in these components. One of the recurrent problems is the frequent failure of weldments due to Type IV cracking (see below), in weldments of ferritic steels subjected to creep loading. Another problem encoun­tered during service exposure of joints of dissimilar ferritic steels is the failure due to the formation of hard brittle zone at the heat-affected zone (HAZ). Both these issues are discussed below.

The modified 9Cr-1Mo steel fusion weld joint (Figure 14) consisting of base metal, deposited weld metal, and the HAZ produces a complex heteroge­neous microstructure due to thermal cycle. The base metal and weld metal consist of a tempered martens­ite structure, with columnar grains in the weld metal.

The HAZ comprises coarse prior-austenitic grain martensite, fine prior-austenitic grain martensite and an intercritical structure, as one traverses from the weld fusion interface toward the unaffected base metal. This is dictated by the peak temperatures experienced by the base metal during weld thermal cycle and the phase transformation characteristics of the steel. It has been established that the localized microstructural degradation in the intercritical region of HAZ is mainly responsible for the prema­ture creep-rupture strength of Cr-Mo weld joint and can be overcome if residual stresses of the weld are adequately relieved by PWHT

The lower creep-rupture strength of weld joint than the base metal is due38,77 to the different types of cracking developed during creep exposure. Four types of cracking have been identified (Figure 15) in Cr-Mo steel weld joint. They have been categor­ized as Type I, Type II, Type III, and Type IV. The Type I and Type II cracks originate in the weld metal, propagate either through the weld metal itself (Type I) or cross over in the HAZ (Type II). The Type III cracking occurs in the coarse grain region of HAZ and can be avoided by refining the grain size. Type IV cracking nucleates and propagates in the intercritical/fine grain region of HAZ. Type IV fail­ure occurs at longer creep exposure and higher test temperature, by coalescence of fine cavities leading to microcracks (Figure 16(a)) and their eventual propagation to the surface.

The type IV cracking susceptibility, defined as the reduction in creep-rupture strength of weld joint compared to its base metal, depends on the type of ferritic steel. 2.25Cr-1Mo steel is most susceptible to type IV cracking; whereas the plain 9Cr-1Mo steel is the least susceptible. At higher test temperature, the type IV cracking susceptibility is higher in modified 9Cr-1Mo steel than the plain steel. This is related77 to the precipitation of Z-phase (Figure 16(b)), a com­plex Cr (V, Nb) N particle, in the modified steel. The Z-phase grows rapidly at elevated temperatures dur­ing long term exposure, by dissolving the beneficial

MX types of precipitates. This promotes the recovery of the substructure with associated decrease in strength in the intercritical region of HAZ.

Although it is difficult to completely eliminate Type IV cracking, several methods are being adopted to improve type IV cracking resistance. It is more severe in thick sections due to the imposed geometri­cal constraint. A design modification can be adopted to decrease the variation in tensile stresses across the welded section of the component or avoid joints in critical regions having high system stresses and relo­cate them in the less critical region. Strength homo­geneity across the weld joint can also be improved by a suitable PWHT. An increase in width of the HAZ can reduce the stress triaxiality such that the soft intercritical region deforms with less constraint with the consequence of reduced creep cavitation, to minimize type IV cracking tendency. The width of the HAZ can be increased both by changing pre­heat and heat-input during welding. Another con­trasting approach to overcome type IV cracking is to avoid or minimize the width of the HAZ, to elimi­nate the intercritical zone. This is being attempted by employing advanced welding techniques such as laser welding. The resistance against intercritical softening can also be improved by increasing the base strength of the steel with the addition of solid solution hard­ening elements such as W, Re, and Co and also by microalloying the steel with boron. Microalloying with boron retards the coarsening rate of M23C6 by replacing some of its carbon. The boron content needs to be optimized with the nitrogen content to avoid BN formation. Addition of Cu is also found to be beneficial. Copper is almost completely insoluble in the iron matrix and when added in small amounts, precipitates as nanosize particles to impart creep resistance. A suitable adjustment of the chemical composition of steel within the specifi­cation range also reduces the large difference in creep strength between the softened HAZ, the base metal, and the coarse grain HAZ of the joint. A weld joint of modified 9Cr-1Mo steel with low carbon, nitrogen, and niobium has been reported to possess creep strength comparable to that of the base steel.

It is expected that a judicious combination of changes in chemistry and process variables would reduce the failures due to type IV cracking in weld­ments of ferritic steels subjected to creep loading.

Another frequent problem78-81 is the formation of ‘hard brittle zone’ during service exposure of dissim­ilar joints between ferritic steels, leading to failures. The formation (Figure 17(a)) of microscopic layer of

hard, brittle zone along the HAZ in dissimilar weld­ments of steels is known to be responsible for the cold cracking, stress corrosion cracking, and higher fre­quency of failures of the weldments. This is one of the cases where modeling has enabled an in-depth understanding of the problem, in addition to providing an industrial solution to prevent the for­mation of brittle zone.

The brittle layer at the interface between 9Cr-1Mo weld and 2.25Cr-1Mo base metal is shown77 to be a manifestation of a number of syner­gistic factors: (a) microstructural changes in regions close to the heat source during welding (b) migration of carbon during PWHT, driven by the gradient in its activity and (c) formation (inset in Figure 17(a))

of series of fine carbides when there is a local supersaturation of carbon. It has been possible to use modeling methods like Finite Difference Meth­ods to predict the carbon profile across the weld region of 9Cr-1Mo and the base metal of 2.25Cr — 1Mo (Figure 17(b)), which were in good agreement with the profiles obtained using electron probe microanalysis. These calculations could be refined using Thermo-Calc and diffusion-controlled trans­formations (DICTRA) to take into account the simultaneous precipitation of carbides and diffusion ofcarbon. These computational methods were instru­mental in predicting the methods to prevent the formation of hard zone in dissimilar joints of ferritic steels. Three elements which would repel carbon

atoms, that is, with the positive interaction energy were chosen for this purpose. Figure 17(c) shows78 the comparison of three different metals, Ni, Cu, and Co in preventing the formation of hard zone. Experimental confirmation was obtained79 using inter­layer between the two dissimilar ferritic steels. Fur­ther insight could also be arrived81 at in the diffusion behavior (Figure 17(d)) of carbon interstitial in the lattices of bcc iron and fcc nickel using molecular dynamics. These calculations could demonstrate that the activation energy for diffusion of carbon in a fcc nickel lattice is higher than bcc iron. This sluggish diffusion kinetics is due to the repulsive potential of nickel toward carbon, which is the main reason for the choice of nickel as the most effective diffusion barrier between the two ferritic steels. Thus, an indus­trial solution to prevent the formation of brittle zone in joints of dissimilar ferritic steels after service exposure could be arrived at, based on an in-depth understanding of the interaction between the lattice potentials of atoms.

It has been demonstrated in the above studies that modeling methods could be used most effectively to reduce the experimental time required for overcom­ing an industrial problem. Experimental benchmark­ing was required only for final confirmation of the predictions. These trends are becoming more common in almost all problems in materials technol­ogy, in recent years, be it atomistic mechanisms or fabrication technologies or prediction of life of com­ponents. It is hoped that this approach of knowledge — based design of materials would gradually replace the time consuming empirical methods of today.

4.03.7 Summary

Future trends in the global fast reactor industry are toward higher operating temperatures, higher burn — up (250 GWdt~ ), higher breeding ratios (~1.4) and longer lifetime for reactor (60-100 years). These goals require several developments in materials sci­ence and technology across all components ofnuclear plants, especially for core component materials.

Ferritic steels have a much better void swelling resistance compared to currently used austenitic stainless steels and are capable of enhancing the burn-up of the fuel up to about ^200GWdt~ Ferritic-martensitic steels based on 9-12% Cr com­positions exhibit the highest swelling resistance and a number of commercial swelling resistant materials have been marketed. The principles behind the design of swelling resistant ferritic steels for core components of fast reactors have been discussed. However, their use is rendered difficult due to their poorer creep strengths at temperatures higher than ~-873 K. Improvement of higher temperature ten­sile and creep strengths in these alloys will enable us to achieve higher temperatures, in addition to higher burn-up, thus improving the economics of nuclear power production. Presently, the reduced creep strength of 9-12Cr ferritic steels at temperatures above 798 K, has restricted their use to certain low stressed components such as subassembly wrappers. Another crucial problem is ‘embrittlement’ in ferritic steels. The mechanisms and methods which are being attempted to overcome embrittlement problems are discussed.

Alloy development programmes are in progress to explore ferritic-martensitic oxide dispersion strength variants, for higher target burn-up of 250 dpa, with enhanced high temperature (^973 K) capability, by improving mechanical properties. Conventional alloy melting routes will have to be abandoned in favor of powder metallurgy techniques of ball-milling, hot isostatic pressing, and hot extrusion for the synthesis of these ODS steels. Process optimization for the development of 9Cr-based ferritic-martensitic steels strengthened by a fine dispersion of yttria nanoparticles has been completed. The major con­cerns in this family of ferritic or ferritic-martensitic steels are the anisotropy of properties in ferritic 12Cr steels or oxidation resistance in 9Cr steels, fabrication procedure, microstructural stability under irradia­tion, and dissolution during back-end technologies.

Materials science, engineering, and technology have become an integral part of the aspiration of the nuclear community to improve the economic viability of fast reactors. One of the major concerns in the alloy development programmes has been the unacceptably long time taken to launch newer mate­rials. It is expected that the current trends in materials development, through intense international collaborations and increased role of modeling in materials behavior, would certainly reduce the time and cost of alloy development programmes for future reactors.