Physical and Welding Metallurgy of Gd-enriched Austenitic Alloys for Spent Nuclear Fuel Applications Part I1: Nickel-based Alloys

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1 WLDING RSARCH SUPPLMNT TO TH WLDING JOURNAL, DCMBR 2004 Sponsored by the American Welding Society and the Welding Research Council Physical and Welding Metallurgy of Gd-enriched Austenitic Alloys for Spent Nuclear Fuel Applications Part I1: Nickel-based Alloys Tests proved Gd-enriched Ni-based alloys are excellent candidates for use in storing spent nuclear fuels J. N. DuPONT, C. V. ROBINO, J. R. MICHAL, R.. MIZIA, AND D. B. WILLIAMS ABSTRACT. The physical and welding metallurgy of gadolinium- (Gd-) enriched Ni-based alloys has been examined using a combination of differential thermal analysis, hot ductility testing, Varestraint testing, and various microstructural characterization techniques. Three different matrix compositions were chosen that were similar to commercial Ni-Cr-Mo base alloys (UNS N06455, N06022, and N06059). A ternary Ni-Cr-Gd alloy was also examined. The Gd level of each alloy was ~2 wt-%. All the alk)ys initiated solidification by formation of primary austenite and terminated solidification by a Liquid * Y + NisGd eutectictype reaction at ~ 1270 C. The solidification temperature ranges of the alloys varied from ~ l(10 to 130 C (depending on alloy composition). This is a substantial reduction compared to the solidification temperature range of Gd-enriched stainless steels (360 to 400 C) that terminate solidification by a peritectic reaction at ~1060 C. The higher-temperature eutectic reaction that occurs in the Ni-based alloys is accompanied by significant improvements in hot ductility and solidification cracking resistance. The results of this research demonstrate that Gd-enriched Ni-based alloys are excellent candidate materials for nuclear criticality control in spent nuclear fuel storage applications that require production J. N. DuPONT is Associate Prqfessor and D B. WILLIAMS is Vice Provost jbr Researeh and a Prqf~sop; Department qf'materials Science & ngineering, Lehigh University, Bethlehem, Pa. C. V. ROBINO is with the Technical St~{ff; Joining and Coating Department, andj. R. MICHAL is with the Technical Staff Materials Characterization Dept.,Sandia National Laboratories, Albuquerque, N.Mex. R.. MIZIA is ngineering Fellow, ner 3, and ngineering Technology. hlaho National ngineering and nvironmental Laboratory, Idaho Falls, Idaho. and fabrication of large amounts of material through conventional ingot metallurgy and fusion welding techniques. Introduction Part 1 of this research article (Ref. 1) summarized results on development of Gd-enriched stainless steel alloys for nuclear criticality control in spent nuclear fuel storage applications. In that work, it was shown that Gd additions to a 316Ltype matrix leads to the formation of an intermetallic (Fe,Ni,Cr)3Gd phase that produces a very large solidification temperature range (360 to 400 C, depending on Gd concentration) and severely limits the hot ductility and weldability of these alloys to a point where commercial production is not practical. As shown by the binary Fe-Gd phase diagram in Fig. IA (Ref. 2), Fe-Gd alloys with low Gd concentrations exhibit a primary delta solidification mode that is followed by a brief region of austenite solidification. Under nonequilibrium solidification conditions in which solute diffusion in the solid is negligible, austenite formation is followed by a series of cascading peritectic reactions before solidification terminates at 845 C by a terminal KYWORDS Gadolinium-nriched Nickel-Based Alloys Austenitic Alloys Differential Thermal Analysis Hot Ductility Testing Varestraint Testing Solidification Cracking utectic Reaction eutectic reaction involving the Fe2Gd intermetallic. Thus, the solidification temperature range of simple Fe-Gd alloys is also very large under nonequilibrium solidification conditions. In multicomponent Gd-enriched stainless steels, solidification starts with primary delta and terminates by a peritectic reaction involving the (Fe,Ni,Cr)3Gd phase at C, which also produces a very large solidification temperature range. Thus, although there are significant differences between simple Fe-Gd alloys and multicomponent Fe-Ni-Cr-Mo-Gd stainless steels, the alloys are similar in that a low-temperature peritectic reaction is responsible for producing a very large solidification temperature range in each system. As mentioned previously, this severely limits the hot ductility and weldability of these alloys. Comparison of the Ni-Gd (Fig. 1B) and Fe-Gd systems reveals some significant differences in the solidification behavior of alloys with low Gd concentrations. In particular, Ni-Gd alloys with less than about 13 wt-% Gd exhibit a simple two step solidification sequence consisting of primary austenite (Ni) formation followed by a terminal eutectic reaction at 1275 C involving the Ni 17Gd2 intermetallic. The presence of the high-temperature eutectic reaction in the Ni-Gd system significantly decreases the solidification temperature range compared to Fe-Gd alloys. Thus, in general, it appears that the solidification behavior of the multicomponent Gd-enriched stainless steels mimics the Fe-Gd system more closely than the Ni- Gd system. From a technical standpoint, it is highly desirable to identify alloying strategies that could be utilized to modify the solidification behavior of the commercial-type Gd-enriched alloys so that solidification more closely follows that of the WLDING JOURNAL ~]i1~.=1

2 O o A 1600, 1400! ~1200 }. looo 4) F" 1 ~ c 03~C -- g12"c "-- (afo) ~_..~._j_ ' I ' ' ' q ' ' 1~o'c Fe-Gd "._ WLDING 1 RSARCH B 14oo o( ) 4) I-- 6o0 lat~ - '. 1~'c \ L Ni-Gd / /\ _c-t " i ' i ' i ' ' J I 87! (Nil-- i i i I o 20 Fe ~% Gadolinium Gd o Gd Wt% Nickel loo Ni Fig Binary phase diagrams. A -- Fe-Gd; B -- Ni-Gd. Table I -- Chemical Compositions of Small-Scale Alloys Used for Preliminary xperiments (All values in wt-%) Alloy Gd Fe Ni Cr Mo Mn Si Bal Bal <0.5 <0.08 Table 2 -- Compositions (wt-%) for the Trial Ni-based Alloys (a) lement N06455-Gd N06022-Gd N06059-Gd Ni-Cr-Gd Gd Mo Cr Fe W Co <0.I <0.10 C Si Mn V P <0.010 <0.010 <0.010 <0.010 S Ti AI Cu N O Ni bal bal bal bal Austenite Fig SM photomicrograph and BSP analysis of high-ni stainless steel Alloy -1 with Gd addition. (a) Values shown are averages of three determinations at each of three laboratories. Ni-Gd system. In particular, it is of interest to develop alloying strategies that would lead to replacement of the lowtemperature peritectic reaction with a higher temperature terminal eutectic reaction. This could potentially produce a significant reduction in the solidification temperature range and concomitant improvements in weldability and hot ductil- ity. The most obvious approach to accomplish this modification would be to increase the Ni content of the matrix. Thus, the objective of this research is to investigate the use of Ni-based alloys for improving the hot ductility and weldability of Gd-enriched austenitic alloys for spent nuclear fuel applications. xperimental Procedure Preliminary Alloy xperiments The solidification responses of two small-scale (2.3 kg) alloys were first evaluated in the as-cast condition before larger scale heats were prepared. The composi- lcp,,[~l$.~ DCMBR 2004

3 WLDING RSARCH 0.160, / ~ exotherm ] ~ exothermic \ , o.t5o -I o~ o.14o t ~ " \ ) "1051,,, ' J 1127 I 1154 "~ ~ ~ q q I A Temperature, C B Temperature, C Fig D TA traces for as-cast high-ni stainless steel Alloy -1 with Gd addition. A -- Heating trace; B -- cooling trace. Fig Light optical photomicrographs of high-ni stainless steel Alloy -1 with Gd addition after DTA analysis. tions of the two experimental alloys are shown in Table 1. Alloy -1 was based on a stainless steel type composition with high Ni, while Alloy -2 is a Ni-based alloy with a matrix composition similar to the commercial alloy UNS #N Alloy - l was examined to determine if an increased Ni content could be used to significantly modify the solidification behavior in a favorable way, while still maintaining a stainless-steel-type matrix composition. The N06022 alloy heat was examined to determine if higher Ni contents were needed in the matrix in order to produce the desired result. The N06022 matrix composition was chosen because this commercial alloy is already being considered for spent nuclear fuel applications. ach alloy was characterized by differential thermal analysis (DTA) and microstructural characterization techniques as described in Part 1 (Ref. 1). Large-Scale Alloy xperiments Results from the preliminary alloy optimization experiments showed that desirable results were obtained by adding Gd to a N06022-type matrix. Thus, four larger scale heats of Ni-based alloys were prepared for more detailed investigations using the same techniques described in Part 1. The compositions of the four alloys are summarized in Table 2. Three of the alloys were chosen to provide matrix compositions similar to highly corrosionresistant Ni-Cr-Mo alloys (UNS #N06455, UNS #N06022, and UNS #N06059), which will provide the proper long-term corrosion resistance under storage conditions. As with the stainless steel alloys, in order to achieve the desired matrix composition, modifications to the nominal alloy composition were required to account for Ni depletion and Cr enrichment of the matrix due to formation of Gd-rich intermetallics (Ref. 3), and these were based on the measured composition of the intermetallic in the small-scale N06022-Gd trial heat. The fourth alloy is a simplified ternary Ni-Cr-Gd alloy that was included as a basis for comparison. Although previous work (Ref. 1) considered Gd additions up to 6 wt-%, recent experiments performed at the Los Alamos National Laboratory Criticality xperiments Facility (Ref. 4) indicate that, for the most highly enriched spent nuclear fuel and the current repository container design, a Gd level of 2 wt-% should be adequate to meet criticality control needs. Thus, target Gd levels were set at 2 wt-%. The values shown in Table 2 are averages of three determinations each at three independent laboratories (nine total measurements). For the four alloys, the standard deviation in the Gd determinations, expressed as a fraction of the average value, ranged from 4.4 to 10.7% of the average for the nine measurements. In general, values for the other major elements were in reasonable agreement, with a single standard deviation of approximately 5% of the average value for that element. The same experimental techniques utilized in Part 1 (DTA, hot ductility, Varestraint weldability, microstructural characterization) were conducted on the large scale Ni-based alloys with the following exceptions. The Ni-based alloys were melted by vacuum induction heating, cast into 10-cm-diameter, ll.3-kg ingots, homogenized at 1160 C for 16 h, and hot rolled at 1160 C with moderate reductions WLDING JOURNAL ~P.li,=t

4 - Light - PMA - SM - DTA - PMA o.t~ t ol5 I o.t~ I Heatln~ Cooling o.o7! o.o~ I xothermic 0.03 I Temperature, C Fig. 5 - addition. optical photomicrograph of Ni-based Alloy -2 with Gd Fig. 6 - result for the Ni-based Alloy -2 with Gd addition. A B 70 ~ 56,,, i ' ' ' I ' ' ' i, ' ' I ' ' ' Si - - Cr - - Fe - - Ni - - M~ r Gd ~ 28 c O Distance (pro) 100 Fig. 7 - results. results acqtdred /?om Ni-hased ztlh)y l:'-2 with (;d addition. A - photomicrograph showing location of PMA trace; B - (3-6 mm) per pass to 14-mm-thick by 15.2-cm-wide plate. Frequent reheating was used to maintain the rolling temperature near 1150 C. Following rolling, the alloys were annealed in an argon atmosphere at 1150 C for 4 h and quenched with chilled flowing argon. Varestraint weldability tests were conducted on the N06455-Gd and Ni-Cr-Gd plate. The Varestraint tests were conducted on 165 x mm subsize samples with a current, voltage, and travel speed of 100 A, 9 V, and 3 mm/s, respectively. Augmented strain levels of 1.0% and 3.5% were used. Simple autogeneous welds were also made on the N06455-Gd alloy plate using electron beam welding (BW) and gas tungsten arc welding (GTAW). The electron beam welds were made at sharp focus, an accelerating voltage of 100 kv, various beam currents between 6 and 40 ma, and travel speeds ranging from 6 to 25 mm/s. Sharp focus was defined as the focus setting that yielded the maximum visible heating of a tungsten block at the appropriate beam current, voltage, and final lens-to-work distance. This produced welds ranging in penetration from 1.5 to 3.2 mm. The autogeneous GTA weld was made at a voltage of 14 V, a current of 120 A, and a travel speed of 3.4 mm/s. Results Preliminary Alloy Optimization xperiments An SM photomicrograph of the high- Ni stainless steel heat (Alloy -I) in the as-cast condition is shown in Fig. 2 along with BSP patterns of the phases observed in the microstructure. The lack of retained ferrite in the dendrite cores is apparent, as is the absence of the thin (Fe,Ni,Cr)3Gd rim and terminal ferrite constituents around the interdendritic (Ni,Fe)3Gd phase that were observed in 316L-type stainless steels enriched in Gd (Ref. 1). With the increased Ni content, solidification appears to initiate by the formation of austenite dendrites and terminate by a peritectic-like reaction, since Fe3Gd and Ni3Gd both form by peritectic reactions in the Fe-Gd and Ni-Gd systems, and the (Fe,Ni,Cr)3Gd phase forms peritectically in Gd-enriched 316L stainless steel. Differential thermal analysis of the Ni modified heat is shown in Fig. 3. During heating, liquation of the (Ni,Fe)3Gd phase initiates at 1127 C- Fig. 3A. Thus, the liquation temperature of the (Ni,Fe)3Gd phase in this alloy is raised by about 65 C compared to that in the 316Ltype alloys, which liquates at C (Ref. 1). The cooling portion of the DTA trace indicates that, for the cooling rate used for the DTA analysis (5 C/min), solidification terminates with the formation of two constituents. The microstructure of the DTA samples is shown in Fig. 4 and in- l~t,'~.,l~"! DCMBR 2004

5 WLDING RSARCH Fig Backscattered electron SM photomicrographs of Jbur Ni-based alloys with Gd additions. A -- NO6455-Gd; B -- NO6022-Gd; C -- NO6059-Gd; D -- Ni-Cr-Gd. Table 3 -- Compositions (All Values in wt-%) of NisGd-~pe Phase Observed in As-Cast Ingots Alloy Ni Cr Mo Mn Fe W AI Si Gd N06455-Gd N06022-Gd N06059-Gd Ni-Cr-Gd Table 4 -- Summary of On-Heating DTA Results Sample utectic Type L ~y + NisGd Liquidus Average Melting Temperature, C Temperature, C Temperature Range, C N06455-Gd 1290, , N06022-Gd 1272, , N06059-Gd 1265, , Ni-Cr-Gd 1291, , dicates that at least two constituents are associated with the interdendritic regions. Although these constituents have not yet been identified, consideration of the Fe- Gd and Ni-Gd phase diagrams, and the established tendency of these alloys to form Gd-rich intermetallics (Ref. 2), implies that the intermetallic phases are probably based on the Ni3Gd and Ni7Gd 2 structures. The presence of two distinct Gd intermetallic phases was not apparent in either the microstructural analysis or heating DTA response of the alloy in the as-cast condition, so it is clear that the cooling rate through the solidification temperature range is an important factor that affects microstructural development in this alloy. In any case, although the solidification temperature range of the Nimodified alloy has been reduced by -65 C, the solidification temperature range is still almost 300 C and would not be expected to significantly improve the weldability and hot ductility. The as-cast microstructure of Alloy -2 WLDING JOURNAL l~p~l~,.. ~]

6 i NisGd Nickel Austenite "!:: / 3.~0.~ 2.00.~ o ~ 1.0-O2 14OOG 1.~'-02 12~10 C I ~ o 0.~'~' ~--_Gd Oxide -t.o0-02 ~ m " ~ ~ ~0 l~a Temt3~ (C)13~ peralu re 1, 14~ Fig Backscattered diffraction results for Alloy Fig DTA scans from Alloy NO6455-Gd that were typical for all the alloys: A -- Heating; B -- NO6022-Gd that were typical for all the alloys, cooling. ~.,.~4~ ~ 60.S J. I i0 o Ni-Cr-Gd I N06059-Gd 1 N06022-Gd =.;..... Temperature (degrees C) 1300 Fig DTA microstructure of Alloy NO6455-Gd showing primary austenite cells and an intercellular "~/NifGd eutectic-type constituent. Fig Hot ductility results of Ni-based alloys with Gd additions. is shown in Fig. 5. The eutectic constituent in the microstructure is clearly visible, and it is evident that the primary austenite is continuous with the austenite in the eutectic constituent. The results of the differential thermal analysis are shown in Fig. 6. On heating, a single liquation event initiates at approximately 1285 C and melting is complete near 1391 C. On cooling, some undercooling is apparent, with solidification initiating at 1374 C and terminating with the formation of a single constituent at 1255 C. The single terminal solidification peak is consistent with the LOM photomicrograph shown in Fig. 5 in which a single eutectic-like constituent was observed. Figure 7 shows the results of an electron probe microanalysis (PMA) scan conducted across the cellular substructure of this alloy. The line shown in Fig. 7A denotes the location of the PMA scan. As with the stainless-steel-type alloys, the intermetallic is high in Gd, and there is essentially no Gd dissolved in the austenite matrix. Based on the results presented above, the Ni-based alloy provides the desirable solidification characteristics in which solidification terminates by a high-temperature eutectic-type reaction instead of a low-temperature peritectic reaction. Comparison of the DTA traces with those for the initial 316L-type heats and the Nimodified alloy indicates that the melting temperature range for this Ni-based alloy is significantly smaller, i.e., -100 C for the Ni-based alloy vs C for the stainless steel alloys. Thus, based on these initial results, a full series of experiments was conducted on several commercialtype Ni-based alloys with Gd additions (compositions shown in Table 2). Full-Scale xperiments on Ni-based Alloys Figure 8 shows typical backscattered electron images of the large-scale alloys in the as-cast condition. All of the alloys exhibited a cellular substructure with an intercellular secondary constituent. Figure 9 shows a representative backscattered diffraction result, which shows that the matrix is austenite and the secondary phase within the eutectic constituent is a Ni5Gd-type intermetallic. Gadolinium oxides were also observed. These results were consistent among all the alloys. The compositions of the NifGd-type phases observed in each alloy are summarized in Table 3. The composition of the phase is consistent with the Ni5Gd stoichiometry, with small amounts of dissolved Cr, Mo, Fe, and A1. Figure 10A shows a DTA heating scan and Fig. 10B shows a cooling scan from Alloy N06455-Gd that was typical for all the alloys. On heating, the alloy exhibits an endothermic peak at 1290 C associated with liquation of the ~,/NifGd eutectic- ~P,,~I$:'I DCMBR 2004 I

7 A -~--. ~ ~ K" ~1~,'.. ~. - ~....,~m&, ; ",e I 50um.~,. -",~,.._ ~ "At,L' I Fig Typical photomicrographs of hot ductifity samples that failed: A -- Outside of the hot zone; B -- inside the hot zone. B..,. b -ql' e.. Fig Light optical micrographs of rolled and annealed NO6455-Gd plate in the following orientations: A -- Longitudinal; B -- transverse; and C -- rolling plane. Rolling direction is right to left in A and C and plate thickness is vertical in B. longated gray features are the NisGd interrnetallic, and small spherical black features are Gd oxides. Fig Light optical photomicrographs of the autogeneous electron beam weld made on Alloy NO6455-Gd in the following conditions: A --As polished; B -- etched. type constituent, and the austenite matrix is fully molten at 1400 C. On cooling, the primary austenite phase begins to solidify at 1400 C, and the terminal Liquid --* y + NisGd eutectic-type reaction occurs at 1272 C (18 C undercooling). These peaks are consistent with the initial alloy microstructures and DTA microstructures (Fig. 11), which exhibit a primary austenite phase and intercellular y/nisgd eutectic-type constituent. Table 4 summarizes the on-heating DTA data. Results are shown for two separate tests conducted on each alloy, and the reproducibility is always within 4 C. The melting temperature range for the alloys varies between 103 and 131 C, which is a substantial reduction compared to the original Gdenriched stainless steels considered ( C) (Ref. 1). Figure 12 shows hot ductility results. In general, each alloy exhibits reasonably good ductility up to a temperature of 1200 C. The ductility is lost at 1250 C, which is near the liquation temperature of the eutectic constituent. The N06022-Gd alloy generally exhibited the lowest ductility while the Ni-Cr-Gd and N06455-Gd alloys typically exhibited the highest ductility at each test temperature. At the lower temperatures of 900 and 1000 C, the samples often failed outside the hot zone. Samples that failed both within and outside of the hot zone were examined using light optical microscopy to determine the location and mode of failure -- Fig. 13. The samples that failed within the hot zone (Fig. 13B) generally exhibited significant plastic deformation of both the matrix and NisGd intermetallic. In contrast, samples that failed outside the hot zone (Fig. 13A) exhibited little ductility and significant cracking of the intermetallic phase. The intermetallic cracks were always approximately normal to the tensile axis. The microstructure of the hot rolled N06455-Gd plate in the three principal WLDING JOURNAL [cp,,l."l$.l

8 A FZ, -",. --,-r -_.2 t,",# -, -.," PMZ '-i=: "'=,~..., - ~:~, ----~ Ld,. -. #..t.. ~ '. ~ --.- L8.....,., ( J ~.. " ",,t.4 lle., t-~, ~., ~',,.~ " / '., ~' ~, %.--- ' -"" "l s -'*"" : : 1 ~ _.. - ' j ~(,,', J Fig Light optical photomicrographs of an autogeneous GTA weld made on the Alloy NO6455-Gd. The weld exhibits a microstructure similar to that of the ingots (primary austenite cells with intercellular 7/Ni5Gd eutectic-type constituent). plate orientations is shown in Fig. 14, and is representative of that observed in all the alloys. As shown, the gadolinide distribution is substantially changed during rolling. At the hot rolling temperature used (1150 C), the gadolinides appear to be relatively soft and ductile, and this results in a gadolinide morphology that is elongated in the rolling direction and flattened out in the rolling plane. In a manner similar to the as-cast microstructure of Fig. 9, the rolled plate also contains small spherical Gd oxides that are apparently distributed throughout the austenite matrix as well as within the NisGd intermetallic. Figure 15 shows light optical photomicrographs of an autogeneous electron beam weld made on the N06455-Gd alloy in the as-polished and etched conditions. The weld exhibits columnar grains that grow epitaxially from the base metal, and this grain morphology is typically observed in fusion welds. No solidification i!l cracks or other defects! were observed in this or the other electron ~. beam welds produced in this study. Figure 16 shows the structure of an autogeneous GTA weld. The weld exhibits a microstructure similar to that of the ingots (primary austenite cells with intercellular "y/nisgd eutectic-type constituent). The cell spacing and secondary phase is much finer than the ingots due to the higher cooling rates in the weld. A partially melted zone (PMZ) is clearly distinguishable outside of the fusion zone (FZ). This PMZ bounds temperatures between the liquidus at the PMZ/FZ interface and the Liquid ~ 7 + NisGd eutectic-type temperature at the PMZ/HAZ interface. Within this region, the "(/NisGd constituent will liquate, as shown in Fig. 16C. The PMZ is often a region where liquation cracking will occur in alloys with wide solidification temperature ranges. However, the solidification temperature range in this alloy is relatively narrow (110 C) and comparable to other nickel-based alloys that are readily weldable. Thus, liquation cracking is generally not expected except, as described below, under conditions of high restraint or where macrosegregation is persistent. Figure 17 shows light optical photomicrographs of an isolated region of a GTA weld that contained a crack. Inspection of the cracked region at slightly higher magnification (Fig. 17B) shows that a relatively large amount of the 7/NisGd constituent exists at the edge of the crack. The large amount of 7/NisGd in this area can be attributed to macrosegregation from the original ingot, and this form of cracking should be easily avoided when macrosegregation in the original ingot is prevented by a secondary refining step such as vacuum arc remelting. Figure 18 shows the Varestraint hot cracking results for the N06455-Gd and Ni-Cr-Gd alloys. Results are shown for the total and maximum crack length. As a basis for comparison, Fig. 19 shows Varestraint hot cracking results for the stainless steel alloys (Ref. 1). The results for the stainless steel alloys were acquired as a function of Gd concentration at a fixed strain level of 3.5%. The stainless steel samples were 0.25 in. thick while the Nibased samples were in. thick. The Varestraint welding parameters were identical for each alloy system. The weld size produced on the stainless steel samples was similar to the welds produced on the Ni-based samples. Although direct comparisons cannot be made between the results for the stainless steels and Nibased alloys due to differences in sample size, the very large difference in maximum and total crack length values clearly shows the significant level of improvement in weldability for the Ni-based alloys. In terms of maximum crack length (MCL), the stainless steel alloys with comparable Gd levels (1.9 wt-%) exhibited MCL values near 5 mm, which is significantly higher than that of the Ni-based alloys of 1 to 1.2 mm. Similar results were obtained when total crack length (TCL) was used as the cracking susceptibility indicator. The TCL value for the stainless steel alloy with 1.9 wt-% Gd was 50 mm, which is significantly higher than the TCL value of 5.7 to 7.5 mm for the Ni-based alloys. D i s c u s s i o n The results of this research show that the solidification behavior and resultant hot ductility and weldability of Gd- lcp~.l~-! DCMBR 2004 I

9 WLDING RSARCH Segrega~ region high,. in eutectic-type constituent 3 Fig Light optical photomicrographs of an isolated region of the GTA weld in Alloy NO6455-Gd that contained a crack. A 1A B c 1.o ~0.8 ~ 0.6 "~ 0A. Z II M-Cr-Gd e N06 55~ t t A i ~ I,- 2.5 Ni-Cr-Gd N06455-Gd! J 0.2 t.5 # Strain (%) o.s s s 4.0 Strain (%) Fig Varestraint hot cracking results for the Ni-Cr-Gd and NO6455-Gd alloys. Results are shown for the following: A -- Maximum crack length; B -- total crack length. enriched austenitic alloys depends strongly on the matrix composition. In particular, stainless-steel-type matrix compositions form a low-temperature (Fe,Ni,Cr)3Gd-type intermetallic by a peritectic reaction. This undesirable reaction sequence and concomitantly large melting temperature range can be avoided by the use of austenitic alloys with a Nibased matrix. The DTA results from the Ni-based alloys indicated that, in the absence of undercooling, solidification initiates at the liquidus temperature (in the range of C depending on alloy composition) by the formation of primary y-austenite. ssentially no Gd is dissolved in the austenite matrix. Thus, as solidification proceeds, the liquid becomes increasingly enriched in Gd until the Liquid ~ y + NisGd eutectic-type reaction is reached, at which point solidification is terminated by the eutectic reaction. This reaction sequence and temperature range is generally similar to that expected in the binary Ni-Gd system. Simple binary Ni- Gd alloys with less than about 13 wt-% Gd exhibit a similar two-step solidification sequence consisting of primary austenite formation followed by a terminal eutectic involving the Ni]vGd 2 intermetallic at 1275 C (Ref. 2). By comparison, the multicomponent Ni-Cr-Mo-Gd alloys terminate solidification in the range of C by a terminal eutectic-type reaction involving the NisGd intermetallic. Thus, although the secondary phase within the terminal eutectic constituent is different in each case, the terminal reaction temperatures are very similar. The Gleeble hot ductility test results confirm that the reduced melting temperature range provides improved hot workability (relative to the Gd-stainless alloys) at temperatures above C. As shown in Fig. 12, the hot ductility of the alloys over the temperature range of C can be roughly grouped into highest ductility (N06455-Gd and Ni-Cr-Gd), intermediate ductility (N06059-Gd), and lowest ductility (N06022-Gd). Qualitatively, this response can be rationalized in terms of the concentrations of major alloying elements (Cr, Mo, Fe, W) and gadolinide volume fraction (Gd level) (Table 2). The Ni-Cr-Gd alloy has the lowest substitutional alloying element level, while the N06455-Gd alloy has an appreciably lower Gd concentration (and intermetallic volume fraction) than the other alloys. Conversely, the N06022-Gd alloy has both the highest alloying element level and highest intermetallic volume fraction, while the N06059-Gd alloy is intermediate by these measures. At temperatures below 900 C, the fracture process initiates by localized cracking in the brittle NisGd WLDING JOURNAL [<P.M~=t

10 A 6 B 50 g = 5 4 o 3 O 2 x e: :Z 1 (z A O~...I ~- 10 i Gd concentration (wl %) Fig: Varestraint hot cracking results for tile stainless steel alloys. Resuhs are shown as a fimction of Gd concentration at a fixed strain level 0]'3.5%. A -- Maximum crack lengths; B -- total crack lengths. phase, while ductility at the higher temperature ( C) appears to be limited by liquation of the NisGd phase. In any case, all the alloys were successfully reduced to plate by hot rolling at temperatures near 1150 C. For these working conditions, the NisGd intermetallic constituent appears to be relatively soft and ductile, and an elongated pancakelike morphology is developed. This morphology is likely not optimal from a mechanical properties perspective, but no attempt was made in the current investigation to optimize either the hot working procedures or the resultant microstructure. Such efforts are, however, ongoing. The significant improvement in solidification cracking resistance of the Nibased alloys compared to the Gd-modified stainless steel alloys can also be attributed to the large reduction in solidification temperature range. The N Gd and Ni-Cr-Gd alloys exhibit a solidification temperature range of II0 C and 131 C, respectively. In comparison, the Gd-stainless steel alloys exhibited a solidification temperature range of C. For a given set of welding parameters (i.e., temperature gradient), the size of the crack-susceptible two-phase solid + liquid region behind the fully molten weld pool increases as the solidification temperature range increases. Thus, the distance a solidification crack can propagate in the two-phase region also increases, resulting in higher MCL and TCL values (i.e., higher cracking susceptibility). Note that the cracking susceptibility is also dependent on the volume fraction of the terminal liquid, but the comparisons between the Gd-stainless alloys and the Gd- Ni alloys were made at similar Gd levels (and thus similar volume fractions of terminal constituents). Comparisons can also be made to commercial Ni-based alloys in which a history of weldability has been established through practical applications. For example, Alloys IN718 and IN625 tested with equivalent size samples and welding parameters at strain level of 2.5% exhibit MCL values of and 1.2 mm, respectively (Ref. 5). These MCL values are comparable to those observed here, and these commercial alloys are known to be readily weldable in typical applications where the level of restraint is not large. Thus, based on the solidification temperature range and Varestraint data acquired here, the Ni-based alloys are expected to be readily weldable under most applications where the level of restraint is not very high. Preliminary, conformation of this was provided in the electron beam and GTA welds that were generally crack free. The isolated region of cracking observed in one GTA weld was confined to a region where a local increase in the amount of the y/nisgd constituent (relative to the nominal y/nisgd content) existed due to macrosegregation. The y/nisgd constituent was present as liquid just prior to the end of solidification. In places where the liquid exists in large quantities, it can promote solidification cracking by interfering with the formation of solid/solid boundaries across cells and grains. This continuous grain boundary/intercellular liquid film cannot support solidification shrinkage strains at the terminal stages of solidification, and hot tears form as a result. It is considered that solidification cracking in this alloy should not be a major problem when macrosegregation has been reduced by an intermediate processing step such as vacuum arc remelting. It should be noted that ongoing work, which includes more extensive welding trials and weld schedule development, has not encountered any difficulties with HAZ cracking in narrow gap cold wire feed GTA butt joint welds. Work is in progress to evaluate this issue in more detail and will be presented in a future article. Conclusions The influence of Gd additions on the solidification behavior, hot ductility, and weldability of Ni-based alloys has been investigated. The following conclusions can be drawn from this research. 1. The addition of 6 wt-% Gd to a nominal 20Ni-14Cr-2Mo stainless-steel-type alloy results in a primary' austenite solidification mode and formation of a Gd-rich interdendritic constituent at C. The resulting solidification temperature range of this high-ni stainless steel is still rather large (300 C) and not significantly different than that previously observed in Type 316L stainless steel alloys with Gd additions (360 to 400 C). 2. The Ni-based alloys with -2 wt-% Gd initiated solidification by primary austenite and terminated solidification by a Liquid ~ y + NisGd eutectic-type reaction at C. The solidification temperature range of these alloys (100 to 130 C) is significantly smaller compared to that of Gd-enriched 316L-type stainless steels that terminate solidification by a peritectic reaction at C. 3. The higher temperature eutectic reaction that occurs in the Ni-based alloys is accompanied by significant improvements in hot ductility and solidification cracking! ~',,t:~i DCMBR 2004

11 resistance. These alloys therefore show considerable potential in terms of primary processing by conventional ingot metallurgy and hot working, and secondary fabrication by fusion welding. A cknowledgments This work was supported by the U.S. Department of nergy, Assistant Secretary for nvironmental Management, under DO Idaho Operations Office Contract No. D-AC07-99IDI3727. This work was performed at Lehigh University, Sandia National Laboratories, and Idaho National ngineering and nvironmental Laboratory through support from the Na- WLDING RSARCH tional Spent Nuclear Fuel Program. Sandia is a multiprogram laboratory operated by Sandia Corporation, a Lockheed Martin company, for the U.S. Department of nergy under Contract D-AC04-94AL8500. References 1. DuPont, J. N., Robino, C. V., Michael, J. R., Mizia, R.., and Williams, D. B Physical and welding metallurgy of Gd-enriched austenitic alloys for spent nuclear fuel applications, Part I: Stainless steel alloys. Welding Jourhal 83(I 1): 289-s to 3011-s. 2. Binaly Alloy Phase Diagrams, Vol Materials Park, Ohio: ASM International. 3. Robino, C. V., DuPont, J. N., Mizia, R... Michael, J. R., Williams, D. B., and Shaber, Development of Gd-enriched alloys for spent nuclear fuel applications -- Part l: Preliminary characterization of small scale Gdenriched stainless steels. Journal ~f Materials l~ineering and Pet~mnanc'e 12(2): Loaiza, D. J., Sanchez, R., Wachs, G., and Mizia, R Critical experiment analysis of a neutron absorbing nickel-chromiummolybdenum-gadolinium alloy being considered for the disposal of spent nuclear fuel.journal ()/'Nuclear Materials Management 32( 1 ). 5. DuPont, J. N., Robino, C. V., and Marder, A. R Solidification and weldability of Nbbearing superalloys, WeMing JounTal 77(10): 417-s to 431-s. Preparation of Manuscripts for Submission to the Welding Journal Research Supplement All authors should address themselves to the following questions when writing papers for submission to the Welding Research Supplement: Why was the work done? What was done? What was found? What is the significance of your results? What are your most important conclusions? With those questions in mind, most authors can logically organize their material along the following lines, using suitable headings and subheadings to divide the paper. 1) Abstract. A concise summary of the major elements of the presentation, not exceeding 200 words, to help the reader decide if the information is for him or her. 2) Introduction. A short statement giving relevant background, purpose, and scope to help orient the reader. Do not duplicate the abstract. 3) xperimental Procedure, Materials, quipment. 4) Results, Discussion. The facts or data obtained and their evaluation. 5) Conclusion. An evaluation and interpretation of your results. Most often, this is what the readers remember. 6) Acknowledgment, References and Appendix. Keep in mind that proper use of terms, abbreviations, and symbols are important considerations in processing a manuscript for publication. For welding terminology, the Welding Journal adheres to AWS A3.0:2001, Standard Welding Terms and Definitions. Papers submitted for consideration in the Welding Research Supplement are required to undergo Peer Review before acceptance for publication. Submit an original and one copy (double-spaced, with 1-in. margins on 8 ½ x l 1-in. or A4 paper) of the manuscript. A manuscript submission form should accompany the manuscript. Tables and figures should be separate from the manuscript copy and only high-quality figures will be published. Figures should be original line art or glossy photos. Special instructions are required if figures are submitted by electronic means. To receive complete instructions and the manuscript submission form, please contact the Peer Review Coordinator, Doreen Kubish, at (305) , ext. 275; FAX ; or write to the American Welding Society, 550 NW LeJeune Rd., Miami, FL WLDING JOURNAL kp~l