The Solidification and Welding Metallurgy of Galling-Resistant Stainless Steels

The autogenous welding response of common galling-resistant stainless steels are evaluated and compared

BY C. V. ROBINO, J. R. MICHAEL AND M. C. MAGUIRE

ABSTRACT. The autogenous welding Introduction these alloys and case studies of welded behavior of two commercial galling-re- fabrication were discussed. Maguire, et sistant austenitic stainless steels, Ni- Conventional austenitic stainless al. (Ref. 2), examined the heat-affected tronic 60 and Gall-Tough, was steels typically display relatively poor zone (HAZ) cracking behavior of Ni- evaluated and compared. The solidifi- galling resistance in many applications, tronic 60 using gas tungsten arc weld- cation behavior and fusion zone hot- and, as a result, a variety of nonstan- ing (GTAW) and HED processing, as cracking tendency of the alloys was dard austenitic steels have been devel- well as weld thermal cycle simulations. evaluated by using differential thermal oped to overcome this difficulty. These Depending on the GTAW conditions, analysis, Varestraint testing and laser steels are normally high-silicon, high- both HAZ liquation and subsolidus spot-welding trials. Gleeble thermal manganese, nitrogen-strengthened al- cracking were observed. For HED over- cycle simulations were used to assess loys and include such grades as lapping spot welds, HAZ cracking was the hot ductility of the alloys during both Nitronic 60, Gall-Tough and Gall- observed under essentially all process- on-heating and on-cooling portions of Tough PLUS. Fusion welding of these ing conditions. weld thermal cycles. Solidification mi- alloys by conventional filler metal pro- There have been a number of investi- crostructures were characterized by cesses such as shielded metal arc (SMA) gations of the weld solidification behav- light optical and electron microscopy, and cold wire feed gas tungsten arc ior, solidification mode and weld and the solidification modes and phases (GTA) welding is normally not prob- cracking behavior of nitrogen-alloyed were identified. Gas tungsten arc (GTA) lematic (Ref. 1). For these processes, and high-manganese, nitrogen-strength- welds in both alloys solidified by the fer- standard filler metals (e.g., E/ER 240 ened stainless steels (Refs. 3-7). In addi- ritic-austenitic mode, and their behavior and E/ER 308) and specially developed tion, there have been several was best described using chromium and fillers (e.g., ER 218) generally provide comprehensive reviews detailing the nickel equivalents developed specifi- satisfactory welding response and many aspects pertaining to the welding cally for the Nitronic series of alloys. weldment performance; however, in sit- of austenitic stainless steels (Refs. 8, 9). Both alloys were found to be somewhat uations requiring autogenous and/or In general, the results of these works in- more susceptible to solidification hot high-energy density (HED) processing, dicate that the effects of manganese and cracking than conventional austenitic the solidification behavior and cracking nitrogen can be rationalized with current stainless steels, although the cracking tendency of the alloys is important. austenitic stainless steel theory, although resistance of Nitronic 60 was somewhat The autogenous weldability of this the roles of these elements are relatively superior to Gall-Tough. Laser spot- class of alloys has not been reported in complex (Ref. 5). welding trials resulted in both fusion the literature, although the general In the case of the three alloys men- and heat-affected zone cracking in the welding behavior of Nitronic 60 was tioned above, the alloy design approach Nitronic 60, while Gall-Tough was re- evaluated by Espy (Ref. 1 ). In that work, appears to be similar, although there are sistant to cracking in these high-solidifi- Espy developed a modified Schaeffler some differences in the balance of alloy cation-rate welds. Comparison of the diagram to describe the ferrite content additions. These differences are most laser weld microstructures indicated of welds in the Nitronic series of alloys. distinct between Nitronic 60 and Gall- that Nitronic 60 shifts to fully austenitic In addition, mechanical properties for Tough, while the Gall-Tough PLUS alloy solidification, while Gall-Tough shifts to is generally similar in composition to an austenitic-ferritic solidification mode Nitronic 60. The Gall-Tough alloy is bal- in high-energy-density processing. The anced to a somewhat higher Creq/Nie~ ratio than Nitronic 60, which implie~ hot ductility measurements indicated KEY WORDS that Gall-Tough is generally superior to that higher as-solidified ferrite contents Nitronic 60 in both on-heating and Galling Resistant are likely and that the autogenous weld- on-cooling tests, apparently as a result GTAW ability of the two alloys may be differ- of differences in grain size and the Hot Ductility ent. Therefore, the purpose of the mechanism of ferrite formation at LBW present study was to evaluate and com- high temperatures. Solidification pare the solidification and autogenous Stainless Steels weldability of Nitronic 60 and Gall- Varestraint Testing Tough alloys. Fusion, HAZ behavior and C V. ROBINO, J. R. MICHAEL and M. C Weldability laser weldability were examined and MAGUIRE are with Sandia National Labora- evaluated. tories, Albuquerque, N.Mex.

446-S I NOVEMBER 1998 Experimental Procedure tility testing performed on a Gleeble also used on selected samples. In this 1500 thermomechanical simulator. This case, samples were polished (or polished The alloys used in this study were test subjects the alloy to a simulated and etched) sections and were generally from commercial heats of Nitronic 60 HAZ thermal cycle, which is interrupted examined in the backscattered electron and Gall-Tough, and the compositions of by rapid fracture during either the on- or secondary electron imaging modes, the heats are shown in Table 1. From heating or on-cooling portion of the ther- respectively, in a JEOL 6400 or Hitachi these heats, samples for differential ther- mal cycle. The thermal cycle used was $4500 SEM operating at 20 kV. Identifi- mal analysis, hot ductility testing, and determined from the Gleeble software cation of microstructural constituents Varestraint testing were machined. All and represents a cycle very close to the was conducted by combinations of AEM, heats were tested in the mill-annealed fusion boundary of a 3.94 kJ/mm weld in electron microprobe analysis (EPMA) condition. Table 1 also shows the com- a 12.7-mm-thick plate. The peak tem- and through the use of backscattered positions of the 304L and 316 stainless perature for these tests was approxi- electron Kikuchi patterns (BEKP) in the steels used for comparison of hot-crack- mately 15°C (27°F) below the SEM (Ref. 13). ing response. nil-strength temperature (NST), although Electron probe microanalysis was per- Differential thermal analysis (DTA) the effect of peak temperature was also formed on a JEOL 8600 electron micro- experiments were performed using a evaluated in several cases. The NST was probe X-ray analyzer operating at 15 kV, Netsch STA 429 thermal analyzer on determined by imposing a tensile load of a spot size of approximately 1 pm and samples of approximately 750 mg mass. approximately 175 N on a hot ductility beam current of 25 nA. Elemental com- Tungsten (>99.99% purity) was used as sample and ramping up the temperature position data was determined by estab- the reference material. Both the refer- at 100°C/s (180°F/s) until sample failure. lished ~(pz) algorithms with integrated Kct ence and the samples were held in high The hot ductility samples were 6.35- X-ray intensities (Ref. 14). BEKP analysis purity alumina crucibles during testing. mm-diameter rods. During hot ductility was conducted on a JEOL 6400 SEM All tests were run in a flowing-helium at- testing, the samples were loaded at a equipped with a custom-made charge mosphere at heating and cooling rates of crosshead velocity of 2 cm/s and ductil- coupled device (CCD) based detector for 0.33°C/s (0.59°F/s). The peak tempera- ity was measured as the reduction of BEKP (Ref. 13). Patterns were collected at ture during testing was approximately area at the fracture surface. an accelerating voltage of 20 kV and a 1500°C (2732°F). Previous experience As-received material, DTA samples beam current of 1 nA and were obtained (Ref. 1O) with this equipment and proce- and welds were examined by light opti- by stopping the beam on a feature of in- dures indicated that a reproducibility of cal metallography. Sample preparation terest and collecting a BEKP by exposing approximately 2°C (3.6°F) in measured for the microstructural examinations in- the CCD camera for 10-20 s. Phase iden- temperatures could be expected. Inter- cluded mounting and polishing through tification was accomplished by auto- pretation of the DTA curves was con- 0.05-pm colloidal silica using standard matic extraction of the important ducted using the convention established metallographic procedures. Samples crystallographic information from the by Maclssac, et al. (Ref. 11 ). were electrolytically etched using a 10% patterns. The crystallographic parameters Susceptibility to fusion-zone hot oxalic acid solution at room temperature are then used in conjunction with quali- cracking was quantified using the longi- and 6 V. Ferrite fraction measurements tative chemical information determined tudinal Varestraint test (Ref. 12). Vare- were conducted using a calibrated mag- by EDS analysis to search a crystallo- straint samples measuring 165 x 25 x negage. Analytical electron microscopy graphic database. Once a candidate 3 mm were fabricated with the long (AEM) was used for further characteriza- match is obtained, the patterns are simu- dimension parallel to the plate rolling tion of selected samples. For these anal- lated and the simulation is compared direction. Autogenous GTAW was used yses, thin foil specimens were prepared with the original pattern. in the tests. The welding parameters were by electrolytic thinning in a 10% HNO 3 100 A direct current electrode negative, in methanol solution at-40°C (-40°F) Results and Discussion 12.5 V and 3.6 mm/s travel speed. Argon and 15 VDC. A Phillips EM30 AEM, op- was used as the shielding gas (20 L/min erating at an accelerating potential of GTA Weld Solidification Mode and Ferrite flow rate) and electrodes were con- 300 kV and equipped with an Oxford In- Content structed of ceriated tungsten. Two sam- struments energy dispersive spectrome- ples were tested at each augmented strain ter (EDS), was used for foil analysis. Figure 1A, B shows the microstruc- over the range from 0.25 to 3.5%. Cracks Scanning electron microscopy (SEM) was tures and BEKP analysis of the various in the fusion zone were measured by using an optical measuring microscope Table 1 -- Alloy Compositions with fluorescent illumination. The pri- mary measures of cracking susceptibility Nitronic 60 Nitronic 60 were taken as the maximum crack length Element (Gleeble) (Varestraint) Gall-Tough 304L 316 (MCL) and total crack length (TCL) on Ni 8.88 8.31 5.17 9.20 10.10 each sample. Cr 16.75 16.45 16.23 18.31 16.30 Laser welding trials were conducted Mn 8.46 8.36 5.50 1.42 1.22 with a Raytheon SS501 400-W average Si 3.96 3.91 3.48 0.64 0.42 power pulsed Nd:YAG laser. Laser spot Mo 0.38 -- -- 0.25 2.05 welds were produced at sharp focus with Cu ------0.46 0.24 a 10-J, 7.2-ms pulse. Laser welds were Co ------0.11 0.23 C 0.078 0.071 0.10 0.021 0.043 produced using the same laser parame- N 0.15 0.15 0.15 0.080 0.051 ters operating at a pulse frequency of 8 Hz P 0.027 0.O29 0.023 0.025 0.031 and travel speed of 1.6 mm/s. S 0.005 0.012 0.020 0.018 0.010 The susceptibility of the alloys to Fe bal. bal. bal. bal. bal. HAZ cracking was assessed by hot duc-

WELDING RESEARCH SUPPLEMENT I 447-s M6X Austenite Ferrite (A) (B)

Fig. I -- Microstructures of GTA welds in Nitronic 60 and Gall-Tough. A -- Nitronic 60, ferrite number 3.9; B -- Gall-Tough, ferrite number I 1.4. Small arrowheads indicate location of eutectic ferrite.

constituents in autogenous GTA welds in was also observed by Ritter, etal. (Ref. 6), Table 2 -- Results of Ferrite Measurements Nitronic 60 and Gall-Tough, respec- in a niobium-modified Nitronic 50-W tively. Both welds were found to solidify weld filler metal that solidified in the fer- Condition Alloy AverageFN in the ferritic-austenitic mode, although ritic-austenitic mode. In the present the quantity of room-temperature ferrite study, the eutectic ferrite was identified Base Metal Nitronic 60 0 differed for the two alloys. For the Ni- by BEKP analysis, shape and location in Gall-Tough 1.5 tronic 60 the average ferrite content of the microstructure and by its tendency to GTA Weld Nitronic 60 3.9 the welds was 3.9 FN, while the average be associated with other minor con- Gall-Tough 11.4 for the Gall-Tough welds was 11.4 FN. In stituents, such as carbides in the case of the ferritic-austenitic solidification mode Nitronic 60 -- Fig. 1A. The BEKPs ob- (Refs. 3, 4, 8, 9, 15), the primary solidifi- tained for the various constituents are first the Welding Research Council (WRC) cation phase is ferrite, while austenite also shown. For both the Nitronic 60 and 1992 diagram (Ref. 8) shown in Fig. 2, it forms in the terminal stage of solidifica- Gall-Tough alloys, the eutectic ferrite is is apparent that this diagram does not ad- tion by a liquid/ferrite/austenite reaction. not a predominant feature of the mi- equately predict either the solidification The primary a-ferrite then transforms to crostructure, although as discussed by mode or ferrite fraction for either alloy. austenite by a diffusional mechanism. Ritter, et al. (Ref. 6), its presence does in- The WRC 1992 diagram is generally con- Delta ferrite, which is retained at room dicate that the composition of the alloy sidered applicable for manganese con- temperature, is thereby located at the must be close to the eutectic trough. tents up to 10 wt-% (even though it is not cores of the original dendrites. Following The retained ferrite content of the base included in the equivalents used), molyb- Brooks and Thompson (Ref. 9), retained metals and GTA welds are compared in denum contents up to 3 wt-%, nitrogen ferrite formed by this mechanism is re- Table 2. As shown, the wrought Nitronic contents up to 0.2 wt-% and silicon con- ferred to as "skeletal ferrite" in the current 60 alloy is fully austenitic, while the Gall- tents up to 1 wt-%. Of these elements, the work. In addition to the skeletal ferrite, Tough alloy contains a small fraction of silicon content of the present alloys at both alloys also appeared to contain a retained ferrite in this condition. For the 3.5- 4.0 wt-% is significantly out of this small fraction of eutectic ferrite. As has GTA welds, the heat of Nitronic 60 eval- range; therefore, it seems likely that it is a been discussed by others (Refs. 6, 15, uated in this work has a ferrite fraction major factor affecting the accuracy of the 16), welds that solidify in the ferritic- close to the lower limit considered ac- diagram for these types of alloys. Elemen- austenitic solidification mode may con- ceptable for avoiding hot cracking. Con- tal segregation patterns determined by tain some eutectic ferrite in addition to versely, the Gall-Tough heat is close to EPMA were essentially similar to those the primary ~3-ferrite. This eutectic ferrite the upper limit of the ferrite range usually observed by Suutala, etal. (Ref. 15), for al- can be distinguished (Ref. 7) from the pri- specified to minimize solidification hot loys solidifying in the ferritic-austenitic mary 6-ferrite by a variety of features, in- cracking. In any case, it is clear that the mode. Silicon generally segregated in a cluding location in the microstructure, ferrite-forming potential of the two alloys manner similar to that for chromium and morphology, solute distribution in the is significantly different. was consequently enriched (to levels ap- austenite adjacent to the ferrite, solute It is notable to compare the ferrite con- proaching 5 wt-%) in the skeletal ferrite. distribution in the ferrite itself, average tent of the GTA welds with that predicted The effects of manganese and n,,itro- ferrite composition and the amount of in- by the various ferrite fraction diagrams gen on the solidification mode and fer- terfacial precipitation. Eutectic ferrite available in the literature. Considering rite content of austenitic stainless steels

448-s I NOVEMBER 1998 are also compli- developed by Hull (Ref. 17) gave the 18 cated. Suutala (Ref. most satisfactory correlation between 5) conducted a composition and solidification mode or 102 4 6 8 10 2 comprehensive ferrite content (if consistency with other evaluation of many austenitic stainless steels is to be main- GTA welds and tained). The equivalents developed by o.C.~ 4-, ¢kl considered the ap- Hull are given by 8o plicability of the ~+ 14 • FA " various (Ref. 8) Cre~ (Hull) = ~r + 1.21(Mo) + 0.48(Si) chromium and + 0.14(Nb) + 2.20(-13) (1) 12 • nickel equivalents, Z~ v C~ ano Nieq, re- Nieq (Hull) = sr,ecuve~v for pre- Ni + 0.1 l(Mn) - 0.0086(Mn) 2 10 diction of + 24o6(C) + 18.4(N) + 0.44(Cu) (2) solidification mode and ferrite 16 18 20 22 24 26 where elemental fractions are given in content. For man- wt-%. Figure 3 shows the solidification Cr Equivalent (%) (Cr + Mo + 0.7 Nb) ganese contents 5 mode and ferrite content estimates for - 8 wt-%, Suutala Nitronic 60 and Gall-Tough using the Fig. 2 -- WRC 1992 diagram showing compositions of Nitronic 60 and (Ref. 5) found that data and correlations of Suutala (Ref. 5) Ga II- Tough. the equivalents and Hull (Ref. 17). The predictions for so-

20 i i i i r i i i i 20 Austenite ~ / :' Aus en,,e .... ' _ Austenite/Ferrite / • • Austenite/Ferrite c / 18 o Ferrite/Austenite J y= 18 o Ferrite/Austenite j / • /- Ferrite ~, /= • • Ferrite A ..~ AF _ Gall Tough ~ /o AF ._ / • Gall Tough ./1' ~ - Nitronic 60 / " ~16 • Nitronic60 o/"/=-,/ ~ 16 >- c • o o _ rt 03 . LLI 14 I,~14 v o • o o o oo _ o • o o Z oo o o z 12~ 12 F o & ~ • • o / 10 o o

8 f I I I I I ~ I I 8 E o "1 I [ I I I 15 17 19 21 23 25 15 17 19 21 23 25 Cr (Hull) Cr (Espy) eq eq (A) (A)

20 i i i i i ~ i F i 20 i i E i i i i /_ • Gall Tough • Gall Tough 0% 18 • Nitronic60 18 0% • Nitronic 60 ~ 5 o/o./_ 50/0/7 16 16

ffl I.~. 14 g W,~14 Z Z 12 12

10 10

8 I I r I I I I I I , 15 17 19 21 23 25 5 17 19 21 23 25 Cr (Hull) eq Creq (Espy)

(B) (B)

Fig. 3 -- Diagrams constructed using Hull chromium and nickel equiva- Fig. 4 -- Diagrams constructed using Espy chromium and nickel equiv- lents. A -- Solidification mode; B -- ferrite content, alents. A -- Solidification mode; B -- ferrite content.

WELDING RESEARCH SUPPLEMENT[ 449-s by Suutala (Ref. 5), those shown in Fig. 3 1373\ 1397X~ provide the best representation of the Ni- tronic 60 and Gall-Tough results. How- ever, these correlations consider a wide Exotherm~ Exolherm~_ ~//1279 range of alloy types and are not opti- mized with respect to the unique com- Endotherm Endotherm positions of the anti-galling steels. Modifications of the original Schaef- On Heating = i ~ OnHeating ~ ~'~1414 ~1291 fler (Ref. 18) equivalents, which are spe- 1386 1286 cific to the anti-galling alloys, were developed by Espy (Ref. 1) to describe the Nitronic series of alloys. These equiva- 1292 lents have the form Cooling OnCooling Creq (Espy) = Cr + Mo+ 1.5(Si) + 0.5(Nb) + 5(V) + 3(AI) (3) / 1357"~" / ~'1392 1429 Nieq (Espy) = Ni + 30(C) 1371 + 0.87 (for Mn)+ 0.33(Cu) + 30(N - 0.045) (4) (A) (B) Espy (Ref. 1) did not construct a solidi- fication mode diagram for these steels. Fig. 5 -- DTA thermograms. A -- Nitronic 60; B -- Gall-Tough for heating and cooling rates of Such a diagram was constructed by Ritter 0.33°C/s (32.6°F/s). Labeled deviations from local baseline are given in centigrade. and Savage (Ref. 6) who replotted the data of Suutala (Ref. 5) for compositions high in lidification mode as shown in Fig. 3A are (Ref. 8). The correlation seems satisfac- manganese and nitrogen using Espy's essentially correct for both alloys, al- tory given the typical accuracy of ferrite equivalents. Figure 4 shows solidification though the estimate for Nitronic 60 is number predictions and the uncertainty mode and ferrite content diagrams from somewhat ambiguous. In terms of ferrite in ferrite number measurements (Brooks the work of Ritter and Savage and Espy, content, the Hull equivalents underesti- and Lippold [Ref. 8] note that variation respectively, and includes the current mate the measured value by approxi- between predicted and measured values alloy compositions. Agreement between mately 4 FN, assuming that FN is commonly differ by as much as 4-8 FN). the observed and predicted solidification approximately equal to volume percent, It is also important to note that of the var- mode and ferrite content is quite good. as is normally the case below about 8 FN ious equivalent formulations evaluated Since these diagrams were developed

if'~"

),, >'i~;7 h;

Ferrite Austenite Austenite Ferrite

(A) (B)

Fig. 6 -- Microstructures of DTA samples. A -- Nitronic 60; B -- Gall-Tough (small arrowheads indicate location of eutectic ferrite).

450-S I NOVEMBER 1998 specifically for nitrogen-strengthened, be consistent with high-manganese steels, it is not surpris- the higher Creq 1.0 .... I .... I .... I .... I .... I .... I .... I .... ing they provide the best description of /Nieq ratio of this • GallTough • E o Nitronic 60 • • solidification mode and ferrite content alloy. The implica- 304L of Nitronic 60 and Gall-Tough alloys. tions of the .¢ • 316 • Nevertheless, it seems appropriate that differences be- the diagrams in Fig. 4 are preferred over tween Nitronic 60 0.6 the more generalized diagrams shown in and Gall-Tough in _~ • o o o O /~ Figs. 2 and 3. terms of on-heat- ing transformation 0 0.4 Differential Thermal Analysis and melting be- havior are dis- E Figure 5 compares differential ther- cussed below in ;02 /i ...... ,,,, mal analysis (DTA) traces from the Ni- detail. Melting of .,.:+~...~ ...... , ,. tronic 60 and Gall-Tough obtained at 8-ferrite com- 0.0 ' - I ~ heating and cooling rates of 0.33°C/s mences at about 0.5 1 1.5 2 2.5 3 3.5 4 I I,M (0.59°F/s). The temperatures shown on 1370°C (2498°F) I ¢.@ Augmented Strain (%) = m the diagrams are averages from two tests for both alloys and for each alloy, although the temperatures continues until the (A) , -- for each run did not differ by more than liquidus tempera- a few degrees. In general, the melting ture is reached at I I,I,I and solidification behavior of the two al- 1392°C (2538°F) 6.0 .... , .... , .... , .... , .... , .... , .... , .... I loys is similar, although there are several for Nitronic 60 and • GallTough important quantitative differences. Melt- 1429°C (2604°F) 5.0 o Nitronic 60 / , -- EE * 304L / I I,M ing of the Nitronic 60 commences at for Gall-Tough. ~ 316 ~ J ==* 1292°C (2358°F), while the Gall-Tough Thermogravimet- begins to liquate at 1327°C (2421°F). ric analysis during Both alloys display a double-peaked the DTA runs melting endotherm with the initial en- showed little or no o+o+go +o -...+. dotherm being significantly larger in the (<1 mg) sample Nitronic 60. The initial endothermic weight changes. peak is interpreted as liquation of Thus, it does not austenite, although solid-state transfor- appear that there I'~ 1.0 ° -"" mation of austenite to ferrite may also were any signifi- contribute to the endotherm. Ritter, etal. cant changes in - - I I,M (Ref. 7), have examined the high-tem- the sample com- 0.5 1 1.5 2 2.5 3 3.5 4 I ~= perature phase stabilities in Nitronic 50, position (e.g., Augmented Strain (%) a similar alloy, and have shown that for volatilization of heat treatments closer to equilibrium Mn or N) during (B) i~u (i.e., long-time isothermal heat treat- the analysis. ments) the alloy can be close to 100% $- Upon cooling ferrite at temperatures near the solidus. from the melt, the Fig. 7 -- Varestraint test results for Nitronic 60, Gall-Tough and con- Evidence for the solid-state transforma- Nitronic 60 and ventional austenitic stainless steels. A -- Maximum crack length; B -- total crack length. tion at temperatures below the initial li- Gall-Tough are quation temperature could not be very similar in their resolved by DTA in either Nitronic 60 or behavior in that the fraction of M23X6 carbide or nitride pre- Gall-Tough, but transformation un- primary solidification event (which gen- cipitates along some of the 6-ferrite- doubtedly occurs. Further, it is conceiv- erally occurs after some undercooling in austenite interfaces. These precipitates able that the initial endotherm at this the present experiments) is followed by a apparently form as a result of elemental heating rate may also reflect continuing single terminal solidification reaction. partitioning during the solid-state trans- solid-state transformation to 8-ferrite, Figure 6 shows the room-temperature mi- formation from 8-ferriteto austenite at although the major contribution to the crostructures of the DTA samples. As temperatures below the solidus. peak is from the austenite liquation. Ini- with the GTA welds, solidification initi- It is also interesting to consider an- tial austenite liquation/solid-state trans- ates with the formation of primary 8-fer- other possible interpretation of the DTA formation continues from 1357 to rite and the large initial exotherm of the solidification traces, particularly in the 1359°C (2443 to 2478°F) for the Ni- DTA traces corresponds to this reaction. case of the Gall-Tough alloy. For the Gall- tronic 60 and Gall-Tough, respectively. As indicated by the smaller secondary Tough, the DTA trace returns to the base- The smaller austenite melting en- exotherm, both alloys terminate solidifi- line prior to initiation of the secondary dotherm in the Gall-Tough implies that a cation with a (presumably peritectic) re- on-cooling exotherm, which can be in- greater fraction of austenite in the Gall- action between the three phases -- terpreted as the completion of solidifica- Tough undergoes a solid-state transfor- liquid, 8-ferrite and austenite. In contrast tion. Therefore, it is conceivable that the mation to 8-ferrite during heating below to the higher-cooling-rate GTA welds, the Gall-Tough DTA samples solidify entirely the solidus. As melting initiates, the Gall- DTA samples did not contain eutectic fer- as 8-ferrite, and the secondary exotherm Tough alloy contains a relatively small rite, but, as shown in Fig. 6B, the Gall- represents the solid-state formation of fraction of austenite, which appears to Tough alloy did contain an appreciable austenite. This view is supported by the

WELDING RESEARCH SUPPLEMENT I 451-s a cooling rate of mented strain. Conversely, compared to * Susceptible 0.33°C/s (0.59°F/s) alloys that are normally considered hot- × Somewhat Susceptible is single phase or cracking susceptible, Nitronic 60 and o Not Susceptible Ga[I-Tough exhibit maximum crack DB Gall Tough Espy not. However, it is Gall Tough Hull important to note lengths -- T~ to 3/~ times, respectively -- [] Gall Tough H&S that since the GTA than that typical of Alloy 718 at an aug- • Nitronic 60 Espy O Nitronic 60 Hull welds in both al- mented strain level of 2.5% under simi- ® Nitronic 60 H&S loys contained eu- lar conditions (Ref. 20). Since the tectic ferrite and Varestraint testing for the conventional 0.20 the DTA samples stainless steels and Alloy 718 was con- Cracking " t No Cracking did not, there are ducted using similar test conditions, some differences travel speeds and weld sizes, compar- FOx 0.15 o~o oo between the solidi- isons between cracking tendencies • • • • o/ o o ¢~ fication mechanics should be reasonable. Thus, autogenous of the DTA and GTA welds in Nitronic 60 and Gall- • oo o lo o ~" 0.10 Xo x GTA weld sam- Tough can best be characterized as + • • o ples. This implies "moderately susceptible to hot cracking." g. • ° ° • o that cooling rate This observation implies that additional • •o So • o o can have a signifi- precautions, in terms of weld joint re- 0.05 • • iv Lo o o x ~ o • x o cant effect on the straint and welding procedures, should °~."*''~_~o~ooOO°° s o solidification be taken when autogenously joining ..w o structure in these these alloys. The results of this work are 0.00 in general agreement with that of Espy 1.0 1.2 1.4 1.6 1.8 2.0 alloys, and is dis- cussed further (Ref. 1), who observed that the fabrica- Cr /Ni eq eq below. tion and service weldability, using bal- anced filler metals, of the Nitronic series Solidification Hot of alloys was acceptable. Fig. 8 -- Regions of cracking sensitivity as a function of impurity content Cracking The hot-cracking response, as mea- and Creq/Nieq ratio. sured by the Varestraint tests, is qualita- The hot cracking tively consistent with the solidification susceptibilities of mode and ferrite content diagrams dis- Nitronic 60 and cussed previously. By both these mea- Gall-Tough, in sures of solidification behavior, GTA terms of Varestraint welds in the two alloys should be rea- maximum and sonably resistant to hot cracking and total crack length, this is reflected in the Varestraint tests. are shown in Fig. 7. Unfortunately, there is no currently By both measures, accepted means for making a direct the heat of Nitronic quantitative comparison between the 60 is somewhat solidification mode or ferrite content and less susceptible to Varestraint cracking behavior. Taken hot cracking than together, these three measures of hot- the Gall-Tough. In cracking response indicate that auto- addition, the mini- genous GTA processing is feasible, mum augmented although, as noted above, additional pre- strain required for cautions relative to conventional observable crack- austenitic stainless steels should be ing is higher in the considered. Nitronic 60 (0.5%) It is well known that the sulfur and than that for the phosphorus contents of austenitic stain- Gall-Tough less steels have a strong influence on Fig. 9 -- Cross sections of laser spot welds. A -- Nitronic 60; (0.75%). Both Ni- weld hot-cracking tendency (Reg. 8, 9). B -- Gall-Tough. tronic 60 and Gall- As a result, the general welding response Tough are of the Nitronic 60 and Gall-Tough in this relatively high Creq/Nieq ratio of the alloy somewhat more context is notable. Figure 8 shows a com- (Creq/Nieq[Espy] -- 1.8) and the observa- susceptible to solidification hot cracking mon means (Refs. 21, 22) of portraying tion that the secondary exotherm is small than alloys that are normally considered solidification cracking susceptibility in in relation to the primary solidification cracking resistant. As shown in Fig. 7, terms of impurity levels and the Creq/Nieq exotherm (as it should be for a solid-state when compared to 304L and 316 grades ratio. The Creq/Nie.H value of 1.5 corre- vs. liquid-solid transformation). In addi- with compositions promoting solidifica- sponds to the change m solidification tion, the fraction of ferrite in the Gall- tion in the ferritic-austenitic mode and mode from primary austenite (below 1.5) Tough DTA samples (FN = 20) is quite tested under the same conditions (Ref. to primary ferrite (above 1.5). For com- high. Unfortunately, from the microstruc- 19), Nitronic 60 and Gall-Tough exhibit positions below 1.5, the boundary tural and DTA evidence alone, it is diffi- maximum crack lengths -- roughly two between crack-sensitive and crack-sus- cult to unequivocally determine whether to four times, respectively-- than that for ceptible welds is sensitive to the impurity or not the solidification of Gall-Tough at the conventional grades at 3% aug- content, while welds above 1.5 remain

452-s [ NOVEMBER 1998 crack-resistant even at relatively high im- means for estimating the solidification purity levels. Brooks and Thompson (Ref. temperature range under conditions sim- 9) have reviewed in detail the various ra- ilar to those encountered in GTAW. Dis- tionales that have been proposed to ex- regarding the effects of fluid flow within plain the superior hot-cracking resistance the weld pool, a reasonable measure of of welds that solidify in the primary fer- the solidification temperature range is rite mode. As shown in Fig. 8, the heats given by the difference between the on- of Nitronic 60 and Gall-Tough are essen- heating liquidus and the on-cooling tially within the no-cracking region of the solidus from the DTA tests (assuming the diagram. For comparison, several equiv- DTA tests for both alloys solidify in the fer- alent equations were used for calculating rite-austenite mode). Using this ap- the Creq/Nie.q ratios for the alloys and are proach, the solidification temperature includecl in the diagram. These equations ranges for the Nitronic 60 and Gall-Tough include those from Hammer and Svenn- are 113 and 144°C (203 and 259°F), re- son (Ref. 23) since these were the equiv- spectively. In essence, this temperature alents used to construct the diagram. Of range provides a measure of the width of the available equivalent formulations, the liquid plus the solid two-phase region those of Hull (Ref. 17) and Espy (Ref. 1) that trails the weld pool. Further, the are shown because they provide the most cracks that form in the Varestraint test at appropriate descriptions of solidification high augmented strains are generally con- mode and ferrite content. sidered to be related to the width of this No matter which equivalent formula- region. The observation that the maxi- tion is used, the Nitronic 60 composition mum crack lengths shown in Fig. 7A tend is significantly closer to the boundary for to saturate at an essentially constant value cracking susceptibility; therefore, it is im- tends to support this interpretation. Thus, Fig. 10--Microstructures of fusion zone. A -- portant to consider why the Nitronic 60 it seems likely that the superior cracking Nitronic 60; B -- Gall-Tough. performs better in the Varestraint testing. response of the Nitronic 60 in the Vare- One possibility relates to the impurity straint test is predominantly a conse- content of the alloys (Table 1). Although quence of its smaller solidification the combined phosphorus and sulfur lev- temperature range. crostructures of Nitronic 60 and Gall- els are similar for the two alloys, the dis- Tough. Both alloys were observed to so- tribution of the two elements is different Laser Welding Behavior lidify in the primary austenite mode (i.e., the sulfur level is higher in the Gall- under these conditions. Cracking in the Tough). The diagrams shown in Fig. 8 High-energy-density (HED) welding is Nitronic 60 (not shown in Fig. 10A) was treat the effects of sulfur and phosphorus often required in the fabrication of minia- observed to occur along solidification with equal weight, which may be an ture components and assemblies and is grain boundaries; however, it should be oversimplification. Ogawa (Ref. 24) ex- generally conducted in the autogenous noted that the dark linear features along amined in detail the individual effects of mode. In addition, the rapid solidifica- the solidification grain boundaries of the phosphorus and sulfur on the Varestraint tion velocities and steep thermal gradi- Gall-Tough alloy in Fig. 10B are not fine hot cracking susceptibility on Invar (a pri- ents associated with HED processing of scale cracks. This observation was veri- mary austenite solidifying iron-nickel austenitic stainless steels are known to fied by examining the laser welds in the alloy). In that work it was found that sul- have dramatic effects on the solidifica- as-polished condition. Thus, these fea- fur was significantly more potent than tion mode and solidification cracking tures appear to be preferentially etched phosphorus in promoting cracking. As a tendency (see extensive reviews by regions along the solidification grain result, it was recommended (Ref. 24) that Brooks and Lippold [Ref. 8] and Brooks boundaries (possibly ferrite or some phosphorus and sulfur be kept below and Thompson [Ref. 9]). For these rea- minor constituent such as MnS), which 0.010 and 0.002 wt-%, respectively. In sons, an evaluation of pulsed laser spot are highlighted by the Nomarski inter- the present work, there clearly is insuffi- and overlapping welds was conducted. ference contrast imaging mode. Unfor- cient data to distinguish between the in- Figure 9 shows the results of laser spot tunately, measurement of the ferrite dividual effects of phosphorus and sulfur, welding trials on the two alloys. As content of the laser welds was precluded but it is conceivable that the difference in shown, Nitronic 60 exhibits extensive fu- by the small size of the welds and by the relative levels of these elements con- sion and heat-affected zone (HAZ) crack- formation of ferrite in the weld HAZ, as tributes to the differences in Varestraint ing, while Gall-Tough does not. These will be discussed below. Based on the cracking response. observations were typical of laser spot etching response, it is assumed that the Another important difference between and seam welds made under a variety Nitronic 60 welds contained little or no the two alloys is apparent from the DTA of conditions. The results are also solidification ferrite, while the Gall- results previously discussed. With respect consistent with previous work (Ref. 2) on Tough contained a small fraction of eu- to weld hot cracking, the solidification Nitronic 60 in which cracking was tectic ferrite. temperature range is an important quan- observed in laser welds made under a As noted above, a change in solidifi- tity (Ref. 25) since it describes the tem- wide range of conditions. cation mode from primary ferrite in GTA perature range over which shrinkage Considering first the fusion zone welds to primary austenite in HED welds strains can develop (and which the re- cracking tendency, it is apparent that is a common observation. It has been maining liquid must support). Although under HED processing the Nitronic 60 is suggested that the change is a result of the DTA experiments do not provide a generally more susceptible than Gall- dendrite tip undercooling (Ref. 9) during measure of the equilibrium solidification Tough to solidification hot cracking. the rapid solidification associated with temperature range, they do provide a Figure 10 compares the fusion zone mi- HED processing. Katayama and Matsu-

WELDING RESEARCH SUPPLEMENT [ 453-s 0.10 .... i ' ' ,11 .... i .... i .... [ .... • Gall Tough Espy 0% 5% 10% Suutala • Gall Tough Espy 22 [] Gall ToughHull • • • . [] GallTough Hull [] Gall Tough Schaeffler 0.08 • Nitronic 60 Espy • Nitronic 60 Espy , ' • " ~ , " 20 O Nitronic 60 Hull (~ Nitronic 60 Hull . . o•..~ . " . Nitronic60 Schaeffler, A . " ~,, •x ~20% 18 ~ 0.06 • Pacaryl 16 m Cracking o • " "H-F" 40% + 14 r.~ 0.04 0 • / o 113 • . , - . • o + .o ./ o"F o- Q. /o No Cracking 12 . • ~ /I x ./ u 80% . ....- 002 °" 10

• , - . - . • o ,

8 i P o ° 0.00 i i i l i i i i i i i . , I .... i .... I .... 1.3 1.4 1.5 1.6 1.7 1.8 1 .g 5 20 25 30 35 Cr /Ni Cr eq eq

Fig. 11 -- HED solidification mode showing location of Nitronic 60 and Fig. 12 -- Effects of impurity level and composition on solidification Gall-Tough compositions. cracking susceptibility of pulsed laser welds in austenitic stainless steels•

to be the most appropriate for the Ni- cles further from the weld interface tronic 60 and Gall-Tough, it is apparent shows no evidence of liquation. Al- that the change to primary austenite so- though the DTA results indicate that the lidification for both alloys is consistent l iquation temperature for the Nitronic 60 with the behavior of conventional is somewhat lower than the correspond- austenitic stainless steels. The signifi- ing temperature for the Gall-Tough and cantly lower Nieq of the Gall-Tough alloy, it is difficult to observe liquation in the however, results in a shift to the primary HAZ area of Fig. 13A, it is reasonable to austenite plus eutectic ferrite mode, conclude that liquation in the Nitronic while the Nitronic 60 shifts to fully 60 does not extend more than approxi- austenitic solidification. From a hot- mately 5 [am from the weld interface of cracking perspective, the austenitic-fer- the laser welds. Figure 9A shows that HAZ cracks in the Nitronic 60 extend to ritic solidification mode (although gener- approximately 50-60 [am from the weld ally inferior to primary ferrite interface; therefore, it is clear that the solidification) is superior to fully HAZ cracks must be predominantly sub- austenitic solidification (Refs. 8, 9). Thus, solidus in character. To more fully char- although both alloys shift to a primary acterize the subsolidus cracking austenite solidification mode, the higher response, a series of hot ductility tests Creq/Nieq ratio of the Gall-Tough is suffi- were conducted on the two alloys. cient to retain some hot-cracking resis- tance during HED processing. This point Hot Ductility Testing is further illustrated in Fig. 12 (a modifi- cation of Fig. 8) to reflect the changes in For Nitronic 60, the range of nil- hot-cracking resistance that accompany strength temperatures (NST) at a 100°C/s HED processing (Refs. 8, 27). (180°F/s) heating rate (for three tests) was Fig. 13 -- Microstructures of laser weld inter- As shown in Fig. 9, cracking in the Ni- 1292-1309°C (2358-2388°F) with an face region. A -- Nitronic 60; B -- Gall-Tough tronic 60 laser welds extends well into (arrows indicate partial liquation of austenite average of 1300°C (2372°F). The range near retained ferrite). the weld HAZ. This cracking is shown in for the Gall-Tough alloy was somewhat more detail in Fig. 13. The HAZ cracks larger, 1278-1317°C (2332-2403°F), extend a significant distance into the with an average of 1303°C (2377°F). For base metal and are believed to consist nawa (Ref. 26) constructed a solidifica- the Nitronic 60, the NST is reasonably primarily of subsolidus cracking, al- consistent with the DTA results, which tion mode diagram for HED welding -- though some liquation cracking near the Fig. 11. As in Fig. 8, several formulations indicated initial liquation at 1292°C weld interface is also probable. Evi- (2358°F) while in comparison to the DTA for the chromium and nickel equivalents dence for the occurrence of a subsolidus for the Nitronic 60 and Gall-Tough are measurements for initial liquidation, the component to the cracking can be ob- NST Gall-Tough measurement was shown in Fig. 11. These equivalents in- tained by observation of the extent of 1327°C (2421°F). Peak temperatures of clude those due to Schaeffler (Ref. 18) austenite liquation adjacent to the re- 1275- 1280°C (2327-2336°F) for Ni- (used to construct the original diagram) tained ferrite in the Gall-Tough -- Fig. tronic 60 and 1285-1290°C (2345- as well as the Hull (Ref. 17) and Espy 13B. Apparent liquation of the austenite 2354°F) for Gall-Tough were conse- (Ref. 1) equivalents (shown earlier to best near the weld interface is apparent describe the behavior of the current within 2-3 [am of the weld interface, quently used to allow for the possibility alloys). Considering the Espy equivalents while austenite adjacent to ferrite parti- of a 15°C (27°F) gradient across the

454-s I NOVEMBER 1998 sample thickness for the on-cooling hot 1263-1266°C ductility tests. A limited number of tests (2305-2311 °F) and using different peak temperatures were these are also shown in 100 also conducted. Fig. 14. As with the GallTough Heating Figure 14 shows the combined results higher peak temperature A 80 GallTough Cooling from: of the hot ductility tests, and it is appar- tests, little ductility is ap- 1285-90°C ~ ent that there are several differences be- parent in the Nitronic 60 ,30+.o °c \ tween the two alloys. The nil-ductility at temperatures well ~ e0 temperature (NDT), or temperature at below the peak temper- Nitmnic60 Heating~~ t c which the ductility drops to zero on heat- ature. Unfortunately, in- o 40 ing, was found to occur at approximately sufficient material was "10-, Nitronic60 Coolingfrom: o \ ~ 1261°C (2302°F) for the Nitronic 60. In /1275.80oc - ~ , |~ NST available to completely n- ~1263-66°C --. ~ ", ~1 GT1303 comparison, the NDT for Gall-Tough was characterize the effect of 2o found to occur at approximately 1293°C peak temperature on (2359°F), which is very close to the aver- ductility recovery for ei- age measured NST. On cooling from a ther alloy. However, it is 700 800 900 1000 1100 1200 1300 1400 peak temperature of 1275-1280°C clear that the hot-ductil- Temperature (°C) (2327-2336°F), the ductility recovery ity response of the two temperature (DRT) for Nitronic 60 occurs alloys is markedly differ- between 910 and 970°C (1670 and ent in both on-heating Fig. 14 -- Hot-ductility test results for Nitronic 60 and Gall-Tough. 1778°F). For Gall-Tough heated to tem- and on-cooling tests. In peratures in the range 1285-1290°C addition, it seems appar- (2345-2354°F), the DRT is close to ent that these differences 1280°C (2336°F). The traditional inter- are likely to be reflected pretation (Ref. 28) of hot-ductility test in the marked difference in laser welding data is that the cracking sensitivity is re- response. lated to the extent of nil-ductility in the Given the basic similarity in composi- HAZ and the rate of ductility recovery on tion between the two alloys, the underly- cooling. From this, it is generally consid- ing reasons for the differences in ered that the magnitude of the tempera- hot-ductility behavior are not readily ture range between the NST and DRT is apparent. Fracture surface analysis indi- a principal indicator of cracking ten- cated that for the peak temperatures used dency. By this measure, the cracking sen- in the majority of the tests -- 1275- sitivity of Nitronic 60 is significantly 1280°C (2327-2336°F) for Nitronic 60 higher than that of Gall-Tough. and 1285-1290°C (2345-2354°F) for the In HAZ regions more remote from the Gall-Tough -- there was generally more weld interface, the peak temperatures are incipient melting in the Nitronic 60. This much lower than the NST. For these re- could lead to generally poorer behavior gions, the NDT also provides an impor- for the Nitronic 60 in the on-cooling tant indicator of the relative cracking tests. The higher fraction of melting does tendency, since it describes the distance not, however, fully explain the differ- from the weld interface that, at some time ences observed for the limited testing of during the weld thermal cycle, has es- Nitronic 60 at lower peak temperatures sentially no ductility. By this measure, the and Gall-Tough at higher temperatures. Gall-Tough is also significantly superior Since the alloys were intended to rep- to the Nitronic 60. resent commercial materials, no attempt Because of the wide range of values was made to ensure that the grain size of encountered in the Gall-Tough NST mea- the alloys was the same. Comparison of surements and the observation that the the as-received materials indicated that NST was significantly lower than the ini- the average grain diameter for the Ni- tial liquation temperature determined by tronic 60 was roughly three to four times DTA, several additional tests were con- larger than that for the Gall-Tough. All ducted to verify the response of the alloy other factors being equal, finer grain-size on cooling. For these tests, a higher peak materials are generally more resistant to temperature of 1305°C (2381 °F) was se- HAZ cracking and, therefore, would be lected and these results are also shown in expected to exhibit superior hot ductility. Fig. 14. At this higher peak temperature, The difference in grain size for the two which is at or above the average materials is notable since the annealing measured NST for Gall-Tough, some temperatures are essentially the same- liquation was apparent. However, the 1066°C (1951°F) for Nitronic 60 and on-cooling ductility response is basically 1038-1093°C (1900-1999°F) for Gall- indistinguishable from the DRT obtained Tough. It seems likely that the retained Fig. 15. -- TEM micrographs of Gleeble simu- from the lower peak temperature tests. ferrite may restrict grain growth during lation samples. A -- Nitronic 60: 1278°C (2332°F) peak temperature, 1020°C (1868°F) Conversely, several on-cooling tests were annealing, so that Gall-Tough may tend test temperature; B and C- Gall-Tough: conducted on the Nitronic 60 using to maintain a finer grain size than Ni- 1290°C (2354°F) peak temperature, 1284°C significantly lower peak temperatures of tronic 60. Further, it also seems likely that (2343°F) test temperature.

WELDING RESEARCH SUPPLEMENT I 455-s is possible that de- of the surface scale results in a net deple- 20 ...... i ' ' i ' | formation incom- tion of Mn and Si in the weld metal, but o Austenite the level of this depletion was not quanti- • Austenite/Ferrite a patibilities due to 18 o Ferrite/Austenite /" strength differ- fied. Clearly, the depletion is likely to • Ferrite be dependent on welding conditions and -4- Gall Tough / ences between the 16 ,-e-- Nitronic 60 Io'~"~ austenite and fer- can affect the solidification characteris- rite, or perhaps dif- tics of the welds. Q. / ¢/) ferences in the More importantly, however, metallo- segregation of graphic examination of the GTA welds z impurity species showed that a region of dense intergran- 12 ' preferentially to ular precipitation formed in both alloys o • the austenite/ferrite in a region 1-2 mm from the weld inter- 10 interfaces vs. face. Clearly, this precipitation is char- austenite grain acteristic of a sensitized region that may I I ? --" "I~'= qp41~ll~--l~ '='l"l~"= p -lll i I , boundaries, may have reduced corrosion resistance. Pre- 16 18 20 22 24 26 have contributed to sumably, a postweld heat treatment Cr (Espy) (PWHT) similar to that used for other eq the poor high-tem- perature ductility. sensitization-prone austenitic stainless (A) Complete charac- steels would probably be effective in terization of the fer- mitigating this problem. rite morphologies, Finally, it is important to reiterate that 19 --l-- Gall Tough 0% /- formation mecha- the current study was conducted on a small number of heats. Figure 16 shows 18 Nitronio 60 f 5%/ _- nism and deforma- tion response the solidification mode and ferrite dia- would require met- grams of Fig. 4 with the composition A 16 +sen, i..... 7, fi10 Q. allographic exami- ranges for each alloy superimposed on iJ /" I nation of samples the diagram. From the size of these g treated to a wide ranges, it is apparent that, like other :~ 13 range of peak tem- austenitic stainless-steel grades, a wide peratures and cool- range of response is possible within a 11 ing conditions. specific alloy designation. However, lO Such an evaluation given the fact that specialty alloys are not was, unfortunately, generally produced at the extremes of outside the scope their composition ranges and specific 5 17 19 21 23 25 of this study. Nev- composition and microstructural differ- C req (Espy) ertheless, it seems ences between the two alloys exist it is apparent that the assumed that the results of the present (B) relative hot ductil- study are generally applicable. ity of the two alloys Fig. 16 -- Diagram showing composition ranges for Nitronic 60 and Gall- is dominated by Summary and Conclusions Tough. A -- Solidification mode; B -- ferrite content. differences in the way ferrite evolves The solidification and welding behav- at high tempera- ior of Nitronic 60 and Gall-Tough have the retained ferrite in Gall-Tough would tures. Further, it is also apparent that the been evaluated by differential thermal tend to restrict grain growth during the transformation to ferrite in the HAZ is analysis, Varestraint testing, hot-ductility hot-ductility testing as well (at least until strongly dependent on the details of the testing and microstructural analysis. The liquation of the austenite occurs). In turn, thermal cycle and, thus, the welding pro- results of this work can be summarized as this would be expected to improve the cedures. follows: Autogenous GTA welds in both relative ductility at high temperatures. alloys solidified by the ferritic-austenitic Finally, some initial investigations of Additional Observations mode, although the higher Creq/Nie. the fine-scale m icrostructures of hot-duc- ratio of Gall-Tough resulted in a highe~ tility test samples were conducted. For There are several additional observa- fraction of ferrite in the welds. The solid- these evaluations, AEM thin foils were tions with respect to autogenous welding ification mode and ferrite fraction of both removed from a region adjacent to the of these alloys that should be noted. First, alloys was best described using chro- fractu re surface of zero ductility on-cool- the Varestraint test apparatus used in this mium and nickel equivalents developed ing tests. TEM micrographs obtained work is equipped with a fully enclosed specifically for the Nitronic series of from these evaluations are shown in Fig. welding area that has been successfully alloys. Both alloys were found to be 15. Ferrite (identified by conventional used for Varestraint testing of titanium somewhat more susceptible to solidifica- electron diffraction) in the Nitronic 60 alloys. Despite this very clean envi- tion hot cracking than conventional HAZ simulations was observed to form ronment, both alloys tend to form a austenitic stainless steels, but autoge- thin films along essentially all of the tenacious surface scale (containing nous GTA processing of the alloys is fea- austenite grain boundaries. Conversely, manganese, silicon and oxygen) that can sible. For GTA welds, the Varestraint ferrite in the Gall-Tough appears to grow affect bead control during welding. A hot-cracking resistance of Nitronic 60 out from the preexisting ferrite and does similar scale was formed on the surface of was superior to Gall-Tough, apparently not generally cover the entire boundary the laser spot welds and can be seen in because of the larger solidification tem- surface. In the case of the Nitronic 60, it the micrograph of Fig. 9A. The formation perature range of the Gall-Tough. Laser

456-s I NOVEMBER 1998 spot welding trials resulted in both fusion Journal 54(6): 189-s to 195-s. martensite formation in stainless steels, weld- and HAZ cracking in the Nitronic 60, 4. Bennett, W. S., and Mills, G. S. 1974. ing Journal 52(5): 193-s to 203-s. while Gall-Tough was generally resistant GTA weldability studies on high manganese 18. Schaeffler, A. L. 1949. Constitution di- stainless steel. Welding Journal 53(12): 548-s agram for stainless steel weld metal. Metals to fusion zone and HAZ cracking in these to 553-s. Progress 56:680 and 680B. high-solidification-rate, high-restraint 5. Suutala, N. 1982. Effect of manganese 19. Robino, C. V., Hills, C. R., and Hlava, welds. Comparison of the laser weld mi- and nitrogen on the solidification mode in P. F. 1994. Characterization of solidification crostructures indicated that Nitronic 60 austenitic stainless steel welds. Metallurgical and weldability of Fe-29Ni-17Co alloys. Low shifts to fully austenitic solidification, Transactions 13A: 2121-2130. Thermal Expansion Alloys and Composites, while the higher Cre~Nie~ ratio of Gall- 6. Ritter, A. M., and Savage, W. F. 1986. eds. J. J. Stephens and D. R. Frear. The Metals Solidification and solidification cracking in Tough tends to limit the s~ift in solidifi- Society of AIME, Warrendale, Penn., pp. nitrogen-strengthened austenitic stainless 123-138. cation mode to an austenitic-ferritic steels. Metallurgical Transactions 17A: 20. Robino, C. V., Michael, J. R., and Cies- solidification mode in HED processing. 727- 737. lak, M. J. 1997. Solidification and welding Hot-ductility measurements indicated 7. Ritter, A. M., Henry, M. F., and Savage, metallurgy of thermo-span alloy. Science and that Gall-Tough is generally superior to W. F. 1984. High temperature phase chemistries Technology of Welding and Joining 2: Nitronic 60 in both on-heating and on- and solidification mode prediction in nitrogen- 220-230. strengthened austenitic stainless steels. Metal- cooling tests and has a very small nil- 21. Kujanpaa, V., Suutala, N., Takalo, T., lurgical Transactions 15A: 1339-1351. and Moisio, T. 1979. Correlation between ductility region. The superior hot 8. Brooks, J. A., and Lippold, J. C. 1993. solidification cracking and microstructure in ductility of Gall-Tough apparently results Selection of wrought austenitic stainless austenitic and austenitic-ferritic stainless steel from differences in grain size between steels. ASM Metals Handbook 6: 456-470. welds. Welding Research International 9: the heats tested, as well as differences in 9. Brooks, J. A., and Thompson, A. W. 55-76. the mechanism of ferrite formation at 1991. Microstructural development and solid- 22. Takalo, T., Suutala, N., and Moisio, T. ification cracking susceptibility of austenitic high temperatures. 1979. Austenitic solidification mode in stainless steel welds. International Metals austenitic stainless steel welds. Metallurgical Reviews 36:16-44. Transactions A 10A: 1173-1181. Acknowledgments 10. Cieslak, M. J., Headley, T. J., Kollie, T., 23. Hammar, O., and Svennson, U. 1979. The authors would like to express and Romig, Jr.,A. D. 1988. A melting and Influence of steel composition on segregation solidification study of Alloy 625. Metallurgical and microstructure during solidification of their thanks to the staffat Sandia National Transactions A 19A: 2319-2331. austenitic stainless steels. Solidification and Laboratories who contributed to this 11. Maclssac,D., Shiohara, B. Y., Chu, M. G., Casting of Metals The Metals Society, London, work: Fred Greulich and Alice Kilgo for and Flemings, M. C. 1983. Structure of pp. 401-410. optical metallography, Paul Hlava and undercooled iron and nickel-based alloys. Grain 24. Ogawa, T. 1986. Weldability of Invar Dick Grant for EPMA analysis, Bonnie Refinement in Castingand Welds, eds. G. J. Ab- and its large-diameter pipe. Welding Journal McKenzie for SEM analysis, Tom Chavez baschian and S. A. David. TMS, 65(8): 213-s to 226-s. pp. 87-96. 25. Borland, J. C. 1960. Generalized for preparation of the AEM thin foils and, 12. Savage,W. F., and Lundin, C. D. 1965. theory of super-solidus cracking in welds (and finally, special thanks to Mike Cieslak for The Varestraint test. Welding Journal 44(10): castings). British Welding Journal 7(8): his valuable comments and thoughtful 433-s to 442-s. 508-512. review of the manuscript. Sandia is a 13. Goehner, R. P., and Michael, J. R. 26. Katayama, S., and Matsunawa, A. multiprogram laboratory, supported by 1996. Phase identification in a scanning elec- 1985. Solidification behavior and microstruc- the U.S. Department of Energy under tron microscope using backscattered electron tural characteristics of pulsed and continuous kikuchi patterns. Journal of Research of the laser welded stainless steels. Proceedings of Contract DE-AC04-94AL85000. National Institute of Standards and Technol- the International Conference on Applications ogy 101 : 301-308. of Lasers and Electo-Optics.**** San Fran- References 14. Scott, V. D., and Love, G. 1983. Quan- cisco, Calif., pp. 19-25. titative Electron-Probe Microanalysis. John 27. Lippold, J. C. 1994. Solidification 1. Espy, R. H. 1982. Weldability of nitro- Wiley, New York, pp. 218-232. behavior and cracking susceptibility of gen-strengthened stainless steels. Welding 15. Suutala, N., Takalo, T., and Moisio, T. pulsed-laser welds in austenitic stainless Journal 61 (5): 149-s to 156-s. 1980. Ferritic-austenitic solidification mode in steels. WeldingJournal 73(6): 129-s to 139-s. 2. Maguire, M. 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