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SUPPLEMENT TO THE WELDING JOURNAL, MARCH 2003Sponsored by the American Welding Society and the Welding Research Council

ABSTRACT. The weld solidification andcracking behavior of sulfur-bearing, free-machining austenitic stainless steel was in-vestigated for both gas tungsten arc (GTA)and pulsed laser beam welding processes.The GTA weld solidification was consis-tent with that predicted with existing so-lidification diagrams, and the cracking re-sponse was controlled primarily by thesolidification mode. The solidification be-havior of the pulsed laser welds was com-plex and often contained regions of pri-mary ferrite and primary austenitesolidification, although in all cases thewelds were found to be completely austen-ite at room temperature. Electronbackscattered diffraction (EBSD) patternanalysis indicated the nature of the basemetal at the time of solidification plays aprimary role in initial solidification. Thesolid-state transformation of austenite toferrite at the fusion zone boundary andferrite to austenite on cooling may both bemassive in nature. A range of alloy com-positions that exhibited good resistance tosolidification cracking and was compatiblewith both welding processes was identi-fied. The compositional range is boundedby laser weldability at lower Creq/Nieq ra-tios and by the GTA weldability at higherratios. It was found with both processesthat the limiting ratios were somewhat de-pendent upon sulfur content.

Introduction

The austenitic stainless steels withminor additions of sulfur, selenium, orlead can exhibit superior machinabilitywith improved surface finishes, reductionin burring, and higher cutting rates. These

alloying elements form second-phase par-ticles that act as lubricants and enhancechip removal during the machiningprocess. The most widely used free-ma-chining stainless steel is AISI 303, in whichsulfur is used at levels as high as 0.4 wt-%.This level is an order of magnitude higherthan the maximum allowable sulfur inAISI 304L of the same basic composition.Concerns regarding the effect of thesehigh-sulfur contents on weld fabricationand performance has limited the use ofsulfur-based stainless steels in applica-tions involving welding.

It is recognized sulfur can be extremelydetrimental to weldability, especially inthe formation of solidification hot cracks.However, Varestraint testing of 303 stain-less steel by Lundin et al. (Ref. 1) indi-cated high levels of sulfur can be toleratedwithout cracking if the weld solidifies asprimary ferrite, while little sulfur can betolerated if the weld solidifies as primaryaustenite. Brooks et al. (Ref. 2) showedsimilar behavior in Fe-Ni-Cr ternary alloyswith high levels of both sulfur and phos-phorus. In conventional welding processessuch as GTA and shielded metal arc(SMA), the change in solidification modefrom primary austenite to primary ferriteoccurs at a Creq/Nieq ratio of ~1.4 whenusing the equivalents of the WRC 92 dia-gram (Ref. 3). This ratio is ~1.50–1.55

(Refs. 4–6) when using Hammar andSvennson equivalents (Ref. 7). For com-positions with Creq/Nieq ratios below thesevalues, the welds typically solidify as pri-mary austenite, whereas, above these val-ues, solidification occurs as primary fer-rite. Other factors such as solidificationvelocity can, however, affect the solidifi-cation behavior.

Little data are available on high-en-ergy-density (HED) welds of the free-ma-chining austenitic stainless steels. It isknown that with the high-solidification ve-locities common with these processes, thetransition in solidification mode from pri-mary ferrite to primary austenite occurs athigher Creq/Nieq ratios than those givenabove, and this is due to dendrite tip un-dercooling (Refs. 8–21). Dendrite tip un-dercooling increases with increasing ve-locity for both ferrite and austenite, butthe rate of this increase is greater for fer-rite. Thus, a critical solidification velocitycan exist above which the austenite phasewill solidify at a higher temperature thanthat of ferrite, and it, therefore, becomesthe stable solidification phase. This criti-cal velocity for transition from primaryferrite to primary austenite is dependenton composition (Refs. 13, 16, 20). In amanner similar to the more conventionalwelding processes, the role of the solidifi-cation mode is very important in the weldcracking behavior during HED processing(at least with nominal levels of P and SRefs. 22–24). It must also be recognizedthat the solidification and solid state trans-formation behavior can be considerablydifferent between the HED weldingprocesses and the more conventional GTAand SMA welding processes (Refs. 9–21).

In the manufacturing of engineeringcomponents, it is common to use severalwelding processes. One region of a com-ponent may require a GTA weld and an-other region a laser beam (LB) weld in the

Weld Solidification and Cracking Behaviorof Free-Machining Stainless Steel

A range of alloy compositions exhibiting good resistance to solidification cracking is identified

BY J. A. BROOKS, C. V. ROBINO, T. J. HEADLEY, AND J. R. MICHAEL

J. A. BROOKS is with Sandia National Labora-tories, Livermore, Calif. C. V. ROBINO, T. J.HEADLEY, and J. R. MICHAEL are with San-dia National Laboratories, Albuquerque, N. Mex.

KEY WORDS

Sulfur BearingStainless SteelLaser Beam WeldingGas Tungsten ArcSolidificationFree MachiningCracking

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same material. In other cases, one heat ofmaterial or one material specification maybe used for the manufacturing of a varietyof components. Thus, alloy compositionsthat exhibit good weldability for bothprocesses would be very desirable. A goalof this investigation was to study the effectof sulfur content and solidification behav-ior on the weld cracking susceptibility ofboth GTA and pulsed YAG laser beamwelding (LBW) in 303 type stainless steeland to identify any differences that mayexist between the two processes. Sum-maries of these studies, including weldproperty data, have been reported else-where (Refs. 25, 26)

Experiments

Ten experimental heats were inductionmelted, cast, and hot rolled into strips

~ 7.6 cm wide and 3.2 mmthick. Chromium andnickel contents were ad-justed to provide a rangeof Creq/Nieq ratios. Sulfurcontent was also variedwhile other minor alloyingelements and impuritieswere held constant. Thealloy compositions andheat designations areshown in Table 1. It can beseen that some of the ma-terials are equivalent to304L with high levels ofsulfur, and others are char-acteristic of the 303 free-machining grade with sul-fur additions as high as 0.4wt-%. Table 1 shows theCreq/Nieq ratios calculatedusing the equivalents ofboth Hammar and Svenn-

son and the WRC 92 diagram. However,for subsequent discussion, the Hammarand Svennson equivalents are used, andcompositions are given in wt-%.

Subsize Varestraint testing (Ref. 27)was used to determine the susceptibility tosolidification cracking of the GTA welds.Composite samples were made by cutting2.5-cm-wide strips across the width of thealloy sheets and electron beam (EB) weld-ing these to 304L end tabs. The compositespecimens were 3.2 mm thick, 2.54 cmwide, and 16.5 cm long. Weld parametersused for the Varestraint test were 90 A, 12V, and 3 mm/s travel speed with argonshielding. Duplicate samples were run ataugmented strains over a range of 0.5 to3.6%. Crack lengths were measured at 30Xmagnification on the as-welded samples.

Autogenous circular welds with a di-ameter of 1 cm were used to assess the

cracking susceptibility of pulsed YAGlaser welds. Two weld schedules wereused, one at 20 Hz, 20 W, and 3.4 mm/s (8in./min) using 2.5 J/pulse and the other at30 Hz, 75 W, and 4.6 mm/s (11 in./min)using 3.7 J/pulse. Single-pulse weld crack-ing susceptibility was also conducted usinga Weeter-type test (Ref. 28) with a hole di-ameter of 0.483 mm and hole depths rang-ing from 0.203 to 0.508 mm.

Samples for metallographic examina-tion were polished and etched with oxalicacid, and samples for transmission electronmicroscopy (TEM) analysis were thinnedusing a percloric acid polishing solution.

Electron backscattered diffraction(EBSD) was used to study the crystallo-graphic orientation of individual grains inthe weld metal. Electron backscattereddiffraction was performed in a JEOL 6400or a JEOL 5900LV scanning electron mi-croscope (SEM). Samples for EBSD wereprepared by standard metallographicpractice followed by light electrolytic etch-ing in oxalic acid or electropolishing in aperchloric acid solution. Electronbackscattered diffraction was conductedat an accelerating voltage of 20 kV with a70-deg sample tilt. It is possible to map theorientation of many grains using EBSD.This was accomplished with the Noran In-struments Orkid system. Automated ori-entation mapping with EBSD is done bythe automated acquisition and analysis ofEBSD patterns for each pixel in an image.In this study the samples were scannedwith grids of 200 x 200 pixels.

Results

GTA Welds

The GTA weld Varestraint results areshown plotted as total crack length vs. aug-mented strain in Fig. 1. The crack lengthreported is the average of duplicate testsat each strain. It can be seen there is alarge range in cracking behavior. Heat 2 isseen to be by far the most susceptible tosolidification cracking, whereas Heats 5and 10 are extremely resistant to cracking(exhibited no cracking even at the higheststrain of 3.6%). Table 2 shows the maxi-mum crack length as a function of aug-mented strain for the same set of tests. Atall strain levels, the maximum crack lengthis greatest again for Heat 2 (Creq/Nieqratio of 1.95 and sulfur content of0.039%). It can also be seen for Heats 1, 4,6, 7, and 9 that the maximum crack lengthsat the highest level of strain tested, 3.6%,are approximately the same. However,their cracking behavior at the lower strainlevels is considerably different. The crack-ing response will be related to solidifica-tion mode and microstructure in a latersection.

Fig. 1 — Varestraint results showing total crack length vs. augmentedstrain.

Table 1 — Alloy Compositions (wt-%) and Calculated Creq/Nieq Ratios

Heat No. 1 2 3 4 5 6 7 8 9 10C 0.03 0.03 0.03 0.03 0.03 0.03 0.03 0.03 0.03 0.03Mn 1.48 1.49 1.48 1.49 1.48 1.48 1.47 1.48 1.49 1.47Si 0.62 0.59 0.62 0.61 0.57 0.57 0.59 0.61 0.59 0.61P 0.03 0.03 0.03 0.03 0.03 0.03 0.03 0.03 0.03 0.03S(a) 0.04 0.04 0.04 0.11 0.27 0.42 0.11 0.12 0.27 0.18Cr(a) 17.55 19.18 18.37 18.45 18.47 18.45 16.87 17.78 17.77 18.16Ni(a) 10.51 8.92 9.73 8.60 8.49 8.61 10.08 9.29 9.30 8.83Mo 0.35 0.35 0.35 0.35 0.35 0.35 0.35 0.35 0.35 0.35Cu 0.31 0.30 0.30 0.30 0.30 0.30 0.30 0.31 0.30 0.30N 0.02 0.02 0.02 0.02 0.02 0.02 0.02 0.02 0.02 0.02

Creq/Nieq 1.48 1.89 1.67 1.85 1.87 1.85 1.48 1.67 1.67 1.78(WRC92)Creq/Nieq 1.55 1.95 1.73 1.92 1.94 1.92 1.55 1.74 1.74 1.85(H and S)

(a) AISI allowable composition of 303: Ni: 8–10%, Cr: 17–19%, S: 0.15% min.

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The alloys can be separated into threegroups to describe their solidification be-havior. Alloys with a Creq/Nieq ratio of1.55 (Heats 1 and 7) generally solidified asprimary ferrite (F/A) but did exhibit someregions of primary austenite solidificationwith small amounts of eutectic ferrite(A/F). Figure 2A shows the overall solidi-fication structure and cracking behavior ofHeat 7 (0.11% S). The darker etching re-gions located primarily near the weld cen-terline solidified as primary austenite, andthe lighter etching regions as primary fer-rite. It can be seen cracking occurred inboth regions. The higher magnificationmicrograph in Fig. 2B shows cracks in theprimary austenite solidified region. Thedark etching spherical particles were iden-tified as sulfides and are primarily locatedin the interdendritic regions of the struc-ture. The cracking behavior of Heat 1(with the same Creq/Nieq ratio as Heat 7but a lower sulfur content of 0.039%) isshown in Fig. 3A. This weld generally so-lidified as primary ferrite, but many of thegrain boundary regions solidified asaustenite, as shown at higher magnifica-tion in Fig 3B. This microstructure andcracking behavior was similar in Heats 7and 1 in the regions of F/A solidification.A SEM micrograph of the region ahead ofthe crack in Heat 7 is shown in Fig. 3C.The cracks are present along austenite so-lidification grain boundaries and are asso-ciated with sulfides. The small sphericalsulfides segregated at the cell boundariesand ahead of a crack tip (some of which

have been etched away during metallo-graphic preparation) are clearly visible inFig. 3C. It was found by SEM EDS analy-sis that the sulfides contain Mn and Cr.Some eutectic ferrite is also apparent inFig. 3C. It appears the regions local to thegrain boundaries were sufficiently modi-fied in composition by segregation and/orbackfilling with Ni- and S-enriched liquidduring Varestraint testing to result inaustenitic solidification with eutectic fer-rite (A/F).

The second group of alloys is thosewith a Creq/Nieq ratio in the range of1.73–1.85 (Heats 3, 8, 9, and 10). These al-loys solidified as primary ferrite and ex-hibited a skeletal ferrite morphology withsome regions of lathy ferrite. An exampleof this microstructure is shown in Fig. 4Afor Heat 10 with a Creq/Nieq ratio of~1.85, tested at a strain of 1%. This heatdid not crack at any strain level tested. Themicrograph was taken near the trailingedge of the weld pool where crackingwould be expected in the Varestraint test.The different etching behavior of the grain

boundary regions is a result of highly lo-calized strain and possibly backfilling dur-ing testing. One such region is shown athigher magnification in the SEM image ofFig. 4B. A large concentration of sulfideparticles are evident at the ferrite / austen-ite boundaries, in the solidification cellboundaries, and especially in the grainboundary regions that exhibited the dif-ferent etching behavior.

The third group of alloys is those withCreq/Nieq ratios in the range of 1.92 to1.95. The microstructure of Heat 2(Creq/Nieq ratio ~1.95) is shown in Fig.5A. Much of the microstructure is charac-teristic of welds that solidify completely asferrite and then transform to austenite ata temperature considerably below that offinal solidification (Refs. 9, 29, 30). In thiscase, the ferrite exhibits a more lath-typestructure in which the laths can extendacross a number of solidification cell di-ameters thereby masking the cell bound-aries. The grain boundary crack shown inthe micrograph is typical of alloys in thisgroup and occurred when tested at an aug-

Fig. 2 — A — Solidification and cracking behavior of Heat 7 containing 0.11% S with a Creq/Nieq ratio of 1.55, dark etching regions solidified as primaryaustenite, lighter regions as primary ferrite; B — higher magnification showing cracks in region of primary austenite solidification.

A B

Table 2 — Maximum Crack Length (mm)

Heat No. 1 2 3 4 5 6 7 8 9 10Strain

0.5% 51.0% 0 15 0 0 0 0 13 8 0 02.5% 10 15 0 11 0 10 11 13 0 03.6% 13 18 8 15 0 15 15 8 15 0

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mented strain of 2.5%. At higher magnifi-cation (Fig. 5B), the grain boundary re-gions are more apparent and are seen tocontain little ferrite and a large concen-tration of sulfides (even at the low sulfurlevel of 0.039%). Although Heat 2 has oneof the lowest levels of sulfur, it exhibitedthe most severe cracking of all the heatsstudied. In heats with this high Creq/Nieqratio, but with the higher levels of sulfur,the microstructures appeared to be morecharacteristic of F/A solidification (moresimilar to that in Fig. 4).

Laser Weld Evaluation

The cracking behavior of laser weldsmade with both weld schedules was as-sessed by metallographically examiningtransverse cross sections every 90 deg inthe circular welds. It was found the twoheats with a Creq/Nieq ratio of ~1.55

(Heats 1 and 7) exhib-ited severe cracking.The sulfur contents ofthese heats are 0.04 and0.11%. Heat 9(Creq/Nieq ratio ~1.74,sulfur level = 0.27%)was the only other heatto exhibit cracking inthese tests, and thecracking was much lesssevere than in Heats 1and 7. The other twoheats with the sameCreq/Nieq ratios similarto Heat 9 but lower lev-els of sulfur, 0.04% and0.12%, did not exhibitcracking. Cracking wasnot observed at any sul-fur level in the heatswith the highestCreq/Nieq ratios of 1.85to 1.95.

The cracking behav-ior of all laser welds is

plotted in Fig. 6 for different impurity lev-els (P+S) and Creq/Nieq ratios. Alsoshown are results of the original diagramof Pacary et al. (Refs. 22, 23) for 304Lstainless steel and those of Lienert (Ref.24) for 304L and 316L. Note that the(P+S) levels of their studies are typical ofthe low-carbon grades of austenitic stain-less steels. The lines separating the regionof cracking and no cracking are a compi-lation of all the data. The justification forplotting the results in terms of combinedphosphorus and sulfur level is based moreon convention than on experimental data(Refs. 5, 31–33). It is recognized both ele-ments can form low melting liquids thatcan lead to solidification cracks, but thepotency of each element may well be dif-ferent. It was found in GTA welds of Fe-Ni-Cr ternary alloys that phosphorus wasmore detrimental to solidification crack-ing response than sulfur (Refs. 2, 9). This

difference was related to the nature of theeutectic liquids, i.e., phosphides formedfilms that more readily wet the solidifyingboundaries, whereas sulfides tended toform as spherical droplets. The effect ofthese two elements on both HAZ and so-lidification cracking is still an area of ac-tive research (Ref. 34). Since the goal ofour work was to determine the effect ofsulfur level, the phosphorus level was heldconstant and near the upper extreme oftypical austenitic stainless steel specifica-tions. In this way, our results should rep-resent worst case behavior for a particularlevel of sulfur in Type 303 alloy, but it mustbe recognized phosphorus may be influ-encing the cracking behavior. It can beseen in Fig. 6 that the results of this studyare consistent with those of the other twoinvestigators. Only the heat with aCreq/Neq ratio of 1.74 and a high level of(P+S) would reflect any change to the ear-lier diagrams (Refs. 22, 24). It should alsobe noted not all the heats are clearly de-fined by the line separating cracking fromno cracking. This is not surprising and issimilar to results reported for GTA welds(Refs. 5, 35).

Laser welds were evaluated using lightmicroscopy, scanning electron microscopy(SEM), and TEM techniques. All weldsconsisted of single-phase austenite withno detectable ferrite. Although the laserweld solidification behavior and mi-crostructures were considerably differentfrom the GTA welds, in general they canalso be separated into three groups basedon Creq/Nieq ratio. Alloys with a Creq/Nieqratio of 1.55 (Heats 1 and 7) solidifiedcompletely as austenite. The austenite so-lidification structure and severe crackingbehavior of Heat 7 with 0.11 sulfur isshown in the low-magnification micro-graph in Fig. 7A. The higher magnifica-tion image of Fig. 7B shows the very dis-tinct cellular solidification structuretypical of primary austenite, with weldcracks confined to the austenite grain

Fig. 3 — A — Solidification and cracking behavior of Heat 1 containing0.04% S with a Creq/Nieq ratio of 1.55; B — higher magnification show-ing primary ferrite solidified structure except in region of crack where finalsolidification occurred as austenite; C — SEM image showing Mn andCr containing sulfides with some eutectic ferrite in cell boundaries.

A B

C

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boundaries. It is interesting to note thegrain or subgrain size appears smallerthan is apparent at the lower magnifica-tion in Fig 7A.

A STEM image from a weld in Heat 1,Creq/Nieq ratio of 1.55, which contains0.04% sulfur, is shown in Fig. 8. In this mi-crograph, the cellular solidification struc-ture resulting from microsegregation isalso very evident. However, the degree ofmicrosegregation was small with a mea-sured cell boundary enrichment in Cr andNi of only ~1 wt-%. The particles notedwith the arrows marked with “O” areamorphous Al-containing oxides, and theparticles marked with “P” are Cr-contain-ing phosphides. However, no sulfideswere observed in the cell boundaries ofTEM samples at this sulfur level (0.04%).

The composition in regions of cracking

in Heat 7 was examined using energy-dis-persive spectrometry and spectrum imag-ing. Spectrum imaging is the collection ofa full EDS X-ray spectrum at each pixel inan image. Once the maps are collected, itis possible to form chemical images. ASEM image of a region analyzed is shownin Fig. 9A, while the maps of phosphorusand sulfur are shown in Figs. 9B and 9C. Itis apparent the phase associated with thecrack is enriched in both these elements.

The second group of alloys is thosewith Creq/Nieq ratios of ~1.74. In general,the solidification behavior of these alloyswas different than those with lowerCreq/Nieq ratios, with the exception ofHeat 9, which contained a higher level ofS, 0.27%. The majority of the weld in Heat9 also solidified as primary austenite, asdid the two heats with the Creq/Nieq ratio

of 1.55. An example of the microstructureshowing the characteristic austenite solid-ified cellular dendritic structure is shownin Fig. 10A and at higher magnification inFig. 10B along with several small inter-granular solidification cracks. The finegrain structure is again apparent. TheWeeter test results on this heat were simi-lar to the circular welds and showed it wasrelatively resistant to cracking, although afew small cracks were observed, as shownin Fig. 10C. This sample was etched morelightly than the sample in Fig. 10B andclearly shows the solidification crack alongan austenite grain boundary. Also appar-ent are darker etching grain boundariesthat contain large quantities of sulfides.

The weld microstructures of the othertwo heats with a Creq/Nieq ratio of 1.74 ex-hibited a different appearance. A repre-

A

A

B

B

Fig. 5 — A — Microstructure of Heat 2 typical of GTA welds in heats with Creq/Nieq ratios of 1.92–1.95. Sample tested at 2.5% strain exhibiting a crack anda lathy ferrite morphology; B — higher magnification showing sulfides in grain boundary region.

Fig. 4 — Heat 10 Creq/Nieq = 1.85 tested at 1% strain. A — Primary ferrite solidified structure exhibiting both skeletal and lathy ferrite morphologies charac-teristic of alloys with Creq/Nieq ratios of 1.73–1.85, note: evidence of high strain in grain boundary regions resulting from Varestraint test; B — SEM image show-ing large sulfide particles at solidification grain boundary (A), ferrite/austenite interface (B), and cell boundaries (C).

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sentative microstructure is shown in Fig.11A for Heat 3 with 0.04% sulfur. Thesewelds contain isolated grains (primarily atthe weld interface and interpulse bound-aries) that exhibit the same cellular solid-ification structure as the laser welds in thelow Creq/Nieq ratio heats. However, themajority of the weld is composed of grainsthat exhibit little or no evidence of a so-lidification structure. Such a region isshown at higher magnification in Fig. 11B.It was concluded, as have other workers(Refs. 9, 12, 19, 23, 25), that grains that ex-hibited little segregation solidified as fer-rite while grains or regions that clearlyexhibited microsegregation or well-pronounced solidification cells solidifiedas austenite. It is also believed the ferritesolidified structure massively transformedto austenite during cooling leaving theweld completely austenitic. The rationalefor this interpretation of the solidificationbehavior is discussed in more detail below.

The third group of alloys (withCreq/Nieq ratios of ~1.85 to 1.95) solidi-fied totally as ferrite and again containedno detectable residual ferrite in the opti-cal, SEM, or TEM microscopes. The mi-crostructure of Heat 2 (Creq/Nieq = 1.95,S = 0.04%) is shown at low magnificationin Fig. 12A, and the overlapping laserpulses are clearly visible. The boundary re-gions are shown at higher magnification inFig. 12B along with ferrite stringers in thebase material. The weld microstructure ofan alloy with a similarly high Creq/Nieqratio, but with high S content (Heat 6, S =0.42), is shown in Fig. 13A. The low-mag-nification micrograph appears very similarto that shown in Fig. 12A, except it is muchdarker due to the high concentration ofsulfides. The fusion zone boundary regionis shown at higher magnification in Fig.13B. In this micrograph, the sulfides arevisible as spherical particles in the fusionzone, as large globules at the fusion zone

boundary, and asstringers (darker etch-ing) in the base mater-ial. Laser welds of thealloys with the highCreq/Nieq ratios werealso studied in TEM. Insamples with the lowersulfur levels, 0.04 wt-%,no sulfides were ob-served although somespherical oxides similarto those in Fig. 8 werenoted. In the alloy con-taining the highest sul-fur content, Heat 6,sulfides exhibit a bi-modal distribution ofsubmicron size parti-cles, as shown in theTEM micrograph inFig. 13C. Many of thelarger particles in thesolidified structure are comprised ofamorphous aluminum oxide surroundedby Cr and Mn containing sulfides. Thelargest particles are contained in the so-lidification cell boundaries, althoughmany of the particles are also presentwithin the columnar dendritic structure.The smaller spherical sulfides noted inFig. 13C are more uniformly distributedthroughout the structure. It is likely theseparticles, ~200 angstroms in diameter,precipitated from the solid state duringweld cooling. These smaller sulfides alsocontained primarily Mn and Cr. It is in-teresting that even with this high sulfidecontent, cracks were not observed in theweld structure.

The austenitic stainless steels are com-plicated by the transformation of austen-ite to ferrite at high temperature. Thus, inGTA welds that solidify as ferrite, a regionof the HAZ from which solidification oc-curs can, depending upon alloy composi-

Fig. 6 — Plot of laser weld cracking behavior.

Fig. 7 — A — Laser weld of Heat 7 with a S content of 0.11% and aCreq/Nieq ratio of 1.55; B — higher magnification of A showing distinctcellular-appearing primary austenite solidification structure and inter-granular nature of cracks.

Fig. 8 — STEM image of laser weld in Heat 7with a S content of 0.04% and a Creq/Nieq ratioof 1.55 showing Al-containing oxides O and Cr-containing phosphides P.

A

B

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tion, be ferrite that formed on heating,which then transforms back to austeniteduring cooling. Experimental evidence ofthis transformation has been observedduring Gleeble testing (Ref. 36), liquid tinweld quenching (Ref. 30), as well as in assolidified weld structures (Ref. 35). Theformation of HAZ ferrite has also beenobserved in situ using X-ray diffractiontechniques (Ref. 37). However, initial so-lidification as epitaxial austenite beforethe onset of ferrite solidification can alsooccur (Ref. 21). It is not clear with the highheating rates of pulsed laser welds if, or towhat extent, a transformation to ferrite oc-curs in the higher Creq/Nieq ratio alloysprior to solidification. The crystal struc-ture at the fusion zone boundary at thetime of solidification would have a majoreffect of the overall weld solidification be-havior. To study this in more detail, elec-tron backscattered diffraction (EBSD)analysis (Ref. 38) was used for both phaseidentification and to determine relativecrystallographic orientations.

A number of fusion zone boundary re-gions were studied in Heat 3 (Creq/Nieqratio ~1.74). As shown in Fig. 11, a mix insolidification mode of ferrite and austen-ite was most apparent at this Creq/Nieq

ratio. First, different re-gions were examined,some with grains thatexhibited a highly cellu-lar structure (primaryaustenite) and otherswith grains that exhib-ited little segregationeffects (primary fer-rite). It was found in allcases that both solidi-fied structures were in-deed face-centeredcubic (FCC) with no de-tectable regions of fer-rite. This observation isconsistent with theTEM results. It wasalso found grains ex-hibiting the distinct cel-lular structure exhib-ited the samecrystallographic orien-tation as the grain fromwhich they grew, sup-porting the hypothesis that epitaxialgrowth of austenite from austenite had oc-curred. Moreover, similar behavior wasobserved at both base metal and inter-pulse weld interfaces. An example at an

interpulse boundary where the austenitediffraction patterns are the same in bothregions is shown in Fig. 14. Regions werealso examined where a cellular solidifica-tion region grew for a short distance be-

Fig. 9 — Position tagged spectroscopy results for laser beam weld in Heat 7 showing A — region of crack; B — phosphorus map; C — sulfur map. Constituentassociated with crack is enriched in both elements.

A B C

Fig. 10 — A — Microstructure of laser weld in Heat 9 with a S contentof 0.27% and a Creq/Nieq ratio of 1.74; B — higher magnification of Ashowing small cracks in the primary austenite solidified structure; C —microstructure of Heat 9 Weeter test sample showing crack at austenitegrain boundary and sulfides along boundaries.

A B

C

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fore leading into a more featureless (fer-rite solidified) structure. Such a region isapparent in Fig. 15 in which the cellularstructure again has the same orientationas the base grain. However, the featurelessregion only a few microns away has a dif-ferent orientation, although microstruc-turally it appears to be within the samegrain. The misorientation between the tworegions (locations 5.1 and 5.2) within theweld is appreciable and is approximatelyequivalent to an 18.5-deg rotation about⟨212⟩. This misorientation is greater thanwould be expected for a small angleboundary or staking fault. The apparentgrowth of ferrite directly from the base

grain in which the base grain and solidifiedregion have different austenite orienta-tions is shown in Fig. 16.

In a similar manner, orientation mapswere used to examine a large region of thefusion zone boundary. An example of pri-mary austenite solidification is shown inFig. 17A in which the sample was tilted 70deg to obtain the Kikuchi patterns. Be-cause of epitaxial growth, it is impossiblefrom the map alone to determine the fu-sion zone boundary. It would seem thefiner-appearing grain structure of the pri-mary austenite welds in Fig. 7B is largelysubgrains with only small orientation dif-ferences. The ferrite solidified welds are,

however, much different. Figure 17Bshows the orientation imaging map of aweld in Heat 2 with a Creq/Nieq ratio of1.95. In this map the region of the fusionzone boundary is more apparent due tothe differences in grain orientation. Con-siderably more grains are apparent on thefusion zone side of the boundary with rel-atively few cases where a single substrategrain corresponds to a single fusion zonegrain. The grain size is also much smaller.Thus it appears that in the case of ferritesolidification, the nucleation and growthconditions are more complex, but consis-tent with the differences in microsegrega-tion behavior.

A

A

B

B

Fig. 11 —A — Microstructure of a laser beam weld in Heat 3 with an S content of 0.04% and a Creq/Nieq ratio of 1.74 showing the nature of mixed mode so-lidification with austenite solidification (dark etching grains) generally confined to region of pulse weld boundaries; B — higher magnification of A taken inregion of weld overlap.

Fig. 12 — A — Microstructure of laser weld in Heat 2, Creq/Nieq = 1.95, 0.04% S; B — higher magnification of boundary regions showing primary ferrite so-lidified structure.

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Discussion

GTA Welds

It has been well established in conven-tional welds, such as GTA, that a strongcorrelation exists between solidificationmode and weld cracking (Refs. 5, 9, 35, 39-42). Welds that solidify as primary ferritewith the secondary solidification ofaustenite, and welds that solidify totally asferrite in which the transformation toaustenite starts at a temperature near thesolidus are extremely resistant to solidifi-cation cracking. Welds that solidify asaustenite are very susceptible to cracking,although this susceptibility can be signifi-cantly reduced at low phosphorus and sul-fur impurity levels. The transition in solid-ification mode from primary austenite toprimary ferrite in low-speed GTA weldsoccurs at a Creq/Nieq ratio ~1.5 whenusing Hammar and Svennson equivalents(Ref. 5) (~1.4 when using the WRC dia-gram). Similarly, welds that solidify totallyas ferrite with a lower transformationtemperature to austenite are also suscep-tible to cracking, although this behavior isnot as well documented. This tendency forcracking at high Creq/Nieq ratios has beenreported by Kujampaa et al. (Ref. 35) tooccur at Creq/Nieq ratios above ~ 2.1.

In examining the Varestraint data inFig. 1, it can be seen the heat most sus-ceptible to weld cracking is Heat 2 with thehighest Creq/Nieq ratio, 1.95, and the low-est sulfur content, 0.04 wt-%. The nextmost susceptible heats, 1, 7, and 4, all ex-hibited similar cracking behavior. Heats 1and 7 both have a Creq/Nieq ratio of 1.55and sulfur levels of 0.04 and 0.11%, re-spectively, The third heat of this group,Heat 4, a Creq/Nieq ratio of 1.92 and a sul-fur level of 0.11 wt-%. The two heats thatexhibited no cracking at any strain levelwere Heats 5 and 10 with Creq/Nieq ratiosof 1.94 and 1.85 and sulfur levels of 0.27and 0.18%, respectively. The overall GTAVarestraint cracking behavior is summa-rized in Fig. 18 and plotted as (S+P) con-tent vs. Creq/Nieq ratio. In general, it can

be seen the data can bebound with a lowerlimit somewhat above1.55 and a higher limitat ~ 1.9. However, itwas observed that theheats with the Creq/Nieqratio ~1.55 containedsome regions solidifiedas austenite, primarilynear the weld center-line, e.g., Fig. 3, Heat 7with 0.11% sulfur.These regions were verysusceptible to cracking.The transition in solidi-fication from primaryferrite to primaryaustenite near the weldcenterline has been re-ported previously inboth GTAW (Ref. 43)and EBW processes(Ref. 44), and has beenattributed to differ-ences in dendrite tip ve-locity. In that work itwas thought the solidi-fication velocity washigher near the weldcenterline — suffi-ciently high to exceedthe critical dendrite tip velocity for thechange in mode from primary ferrite toprimary austenite. This type of behaviorwill typically occur at compositions nearthe critical Creq/Nieq ratio of 1.5, and islikely the cause of the localized change insolidification mode observed here. How-ever, it was also observed that in primaryferrite solidification regions of welds withthe Creq/Nieq ratio ~1.55 relatively severecracking also occurred. This is somewhatsurprising, but it could be seen that in thegrain boundary regions of cracking, thelocal composition was sufficiently alteredto promote primary austenite solidifica-tion (Ref. 42). This may be due partly tobackfilling during straining with liquid suf-ficiently enriched with Ni and S to result inlocal primary austenite solidification —

Fig. 3B and 3C. It may be that at this highlevel of S, slightly higher Creq/Nieq ratios,with a concomitant larger fraction of fer-rite solidification, are required to signifi-cantly reduce cracking. It has been ob-served that in Fe-Ni-Cr ternary alloys atCreq/Nieq ratios as high as ~1.6, high lev-els of sulfur tended to promote austenitesolidification and alter residual ferritecontent (Ref. 45). In the alloys studiedhere, a higher Creq/Nieq ratio of ~1.6 isneeded to ensure good cracking resistanceand may also be due to sulfur acting as aNi equivalent.

The upper and lower bounds of thecracking susceptibility region shown inFig. 18 are consistent with the work of Su-utula and Kujanpää (Ref. 35). However,the limits are somewhat tightened, ~1.6 to1.9, compared to their values of 1.5 to 2.1

Fig. 13 — A — Microstructure of laser weld in Heat 6 with Creq/Nieq =1.92 and 0.42% S; B — boundary regions showing primary ferrite solid-ified structure containing high concentration of sulfides, globular sulfidesat HAZ boundary, and sulfide and ferrite stringers in base material; C -—TEM micrograph showing globular sulfides in solidification boundary,spherical Mn and Cr containing sulfides nucleated at amorphous alu-minum oxide particles, and fine uniformly distributed sulfides.

A B

C

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(using Hammar and Svennson equiva-lents). The observed cracking at 1.55 iswell within the data they compiled, espe-cially considering the high levels of S anddifferent testing techniques. The upperlimit of 1.9 vs. 2.1 is significantly differentand may also be related to weld test tech-niques. Since the change in cracking be-havior is not associated simply with anabrupt change in solidification mode, as isthe case of the lower bound of Creq/Nieqratio, it would be expected that the upperlimit might not be as well defined. How-ever, the change in cracking behavior isstill related to solidification and solid-state transformation behavior. It is clearfrom the microstructures, with ferrite ex-tending over a number of cell boundaries(Fig. 5A), that low S welds with Creq/Nieqratio of ~1.92 solidified completely as fer-rite. Moreover, the morphology of the fer-rite suggests the transformation of ferriteoccurred at a temperature considerablybelow the solidus temperature (Refs. 9,29, 30). The rationale summarized (Refs.2, 9) to explain cracking behavior is con-sistent with the GTA weld data shownhere. Solidification crack initiation andpropagation along the complex interphaseboundaries, formed during either ferrite-austenite solidification, or during thehigh-temperature transformation of pri-mary ferrite to austenite, is more difficultthan along the rather smooth, single-phase grain boundaries (Ref. 46). Thesesingle-phase boundaries occur at both thelow Creq/Nieq ratios with austenite solidi-

fication and at highCreq/Nieq ratioswhere single-phaseferrite solidifica-tion predominates.Also, crack propa-gation may be moredifficult along thelower surface en-ergy δ–γ bound-aries than along thehigher energy δ–δand γ–γ boundariesthat are formed during solidification andcooling (Ref. 47).

The upper bound Creq/Nieq ratio of~1.9 for favorable cracking resistance alsoappears to be dependent upon sulfur con-tent; i.e., the higher the sulfur level at aCreq/Nieq ratio of ~1.93, the greater theresistance to cracking. This behavior wasalso observed in ternary heats doped withhigh levels of sulfur (Ref. 2) and may bethe result of backfilling from sulfur-con-taining eutectic liquid. In Fig. 3, it appearsfrom the etching behavior the grainboundary regions (which did not formcracks) underwent considerable strain andcontain large amounts of sulfides. Theseobservations are consistent with a back-filling mechanism. However, it appearedmicrostructurally that higher levels of sul-fur extended the range of two-phase F/Asolidification to higher Creq/Nieq ratioswhich would also improve cracking sus-ceptibility. This is consistent with S alter-ing the solidification mode at the lower ra-

tios of ~1.55 (Ref. 26).It has been suggested the maximum

crack length at the saturation strain levelmay be a better indicator of cracking be-havior than the total crack length plottedagainst augmented strain. When analyzingVarestraint data, the saturation cracklength at the highest strains correspondsto the temperature-sensitive region forcracking. The length of this region in agiven weld is then directly related to thethermal gradient along the weld center-line (for a centerline crack) and the weldspeed. The highest strain tested, 3.6%,may be close to the saturation strain. Ifthis is the case, the maximum cracklengths at 3.6% strain in Table 2 wouldsuggest Heat 2 is the most susceptible heatand this agrees with the total crack lengthdata of Fig. 1. However, based on maxi-mum crack length, the cracking behaviorof Heats 4, 6, 7, and 9 should all be verysimilar and likely not that different fromHeats 1 and 2. We believe the results inFigs. 1 and 18, rather than the maximum

Fig. 14 — SEM micrograph with EBSD patterns showing epitaxial solidifica-tion of austenite at interpulse boundary. The EBSD patterns show both re-gions have same austenite orientation. (Heat 3, Creq/Nieq = 1.73).

Fig. 15 — SEM micrograph showing cellular structure growing a short dis-tance from the base material and changing to apparent featureless ferrite so-lidification. All EBSD patterns are FCC austenite, and show epitaxial solidi-fication of austenite from base material with same orientation, while theorientation of the region that solidified as ferrite differs from the orientationof the base metal. (Heat 3).

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crack length at 3.6% strain, may be morerepresentative of cracking response inproduction environments for the heats ofmaterial studied here.

Laser Beam Welds

Before discussing the solidificationcracking behavior of the laser beam welds,or correlating solidification cracking be-havior to solidification mode, one must beable to distinguish between the two pri-mary modes of solidification. No weld fer-rite was present that could more clearlydefine solidification behavior. We con-cluded the weld regions that exhibited thedistinct cellular dendritic structure solidi-fied as primary austenite while those re-gions that exhibited little or no evidence ofthe cellular solidification structure solidi-fied as ferrite. These conclusions werebased on earlier studies that showed thecombined effect of dendrite tip under-cooling and solid-state diffusion is muchmore effective in reducing microsegrega-tion during ferrite solidification than inaustenite solidification (Refs. 14, 15).However, other workers have proposedthat the segregation-free, high-energy-density laser welds may be the result ofpartitionless solidification (Ref. 13). In ei-ther case, the use of EBSD further sup-ported the relationship between the pri-

mary phase of solidification and the de-gree of microsegregation

There appeared to be three differenttypes of initial solidification behavior ofthe laser welds, similar to those reportedby Lippold (Ref. 23). The first was the ini-tial epitaxial growth and continuing solid-ification of austenite. The second was theinitial epitaxial growth of austenite fol-lowed by a transition to ferrite solidifica-tion within a distance of only a few mi-crons. Both behaviors existed with aCreq/Nieq ratio of 1.74 at the lower sulfurlevels. The difference in the two behaviorsmay be related to the crystallographic ori-entation of the substrate grain and the di-rection of maximum heat flow. If the indi-vidual grains solidify at a rate near that ofthe macroscopic solidification front, thendendrites within different grains mustgrow at different velocities. Dendrites ingrains with a <100> direction alignedalong that of maximum heat flow (normalto the macroscopic solid/liquid interface)will grow at a slower rate than dendriteswithin grains with significant misalign-ment. Dendrites in these less favorablyaligned grains must travel further, andthus faster, to maintain the velocity of themacroscopic solidification front. How-ever, their tips will also be at a lower tem-perature and lag those in an adjacent,more favorably oriented, grain (Ref. 48).

These grains may then be cut off duringthe commonly observed competitivegrowth process. In stainless steel laserwelds, when the macroscopic interface ve-locity is near that of the critical transitionvelocity for δ→γ solidification (Creq/Nieqratio ~1.7 in this study), the increaseddendrite growth velocity in less favorablyaligned grains may promote austenite so-lidification. When the growth direction ismore favorably aligned and dendrite tipvelocities are slower, a transition from ini-tial solidification of austenite to thegrowth of ferrite will occur. However,growth of ferrite will first require nucle-ation. As a result, the transition to ferritesolidification will occur for velocitiessomewhat slower than that of the com-monly calculated critical transition veloc-ity (Ref. 15).

The third condition appeared to be di-rect solidification of ferrite from the sub-strate. This was the common mode at thehigher Creq/Nieq ratios ~1.9. There ap-peared to be no precursor solidification ofaustenite, and the austenite orientationswere usually different across the interface.It also appeared in many cases that a sin-gle grain of ferrite solidified from a singlesubstrate grain — Fig. 16. This suggeststhe fusion zone boundary region was fer-rite at the time of initial solidification, pos-sibly as a result of a massive transforma-

Fig. 16 — SEM micrograph with EBSD patterns showing solidification as fer-rite from base grain. The EBSD patterns show different austenite orientationsfor the two regions. (Heat 3).

Fig. 17 — A — Orientation imaging map of primary austenite solidified laserweld of Heat 1. Note fusion zone boundary in map is not apparent due to epi-taxial growth; B — similar map of Heat 2. Note orientation change acrossmuch of fusion zone boundary.

A

B

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tion on heating. If ferrite nucleated at theaustenite interface during initial solidifi-cation, numerous grains may coincidewith individual substrate grains unless fer-rite nucleated on austenite with a pre-ferred, low-energy configuration. Low-energy orientations between ferrite andaustenite established during both solidifi-cation (Refs. 49, 50) and solid-state trans-formations (Ref. 9) have been observed instainless steel welds. The EBSD patternsin Fig. 17B covering a large region of theweld show the FZ boundary region con-sists of more grains than that of the corre-sponding substrate. However, this is pri-marily a result of the solidified ferritemassively transforming to austenite, yield-ing a finer grain structure than that of theparent ferrite grain, which complicates in-terpretation of the solidification behavior(Ref. 51). It is also possible the massivetransformation could initiate at theaustenite fusion and interpass boundaries,resulting in no orientation change acrossthe boundary; this was not commonly ob-served. It is evident initial solidification offerrite in the pulsed laser welds is complexand needs further study. Nevertheless, theuse of segregation patterns to distinguishsolidification mode, which is important forthis study, is consistent with the crystallo-graphic orientations observed for the dif-ferent structures.

Irrespective of the details of the solidi-fication mechanics, it was found onlywelds that solidified as primary austeniteexhibited weld cracking. This behavior isconsistent with earlier results on pulsedlaser welds with lower levels of phospho-rus and sulfur (Refs. 22–24). Our data isalso consistent with previous work that in-dicates the change in solidification mode

due to dendrite tip undercooling occurs ata Creq/Nieq ratio of about 1.68–1.7. Theline in Fig. 6 separating cracking from nocracking for Creq/Nieq ratios below ~1.7was defined by the data of Pacary et al. andLienert. In the Creq/Nieq ratio range be-tween ~1.6 and 1.7, their welds often ex-hibited mixed modes of solidification offerrite and austenite. Pacary et al. showeda vertical line separating cracking and nocracking at ~1.67 and Lienert at 1.7.Those lines generally corresponded to thetransition to nearly complete ferrite solid-ification. However, we found that with aCreq/Nieq ratio of 1.74, the welds still so-lidified in a mixed mode, although thefraction of primary ferrite solidificationwas typically high, ~75% or greater. Anexception was Heat 9 with the highestamount of sulfur (0.27%), in which a largefraction of the weld solidified as primaryaustenite. In this heat, however, only asmall amount of cracking was observed. Inthis heat, the solidified grain boundariescontained large amounts of sulfides, whichimplies sufficient eutectic liquid was avail-able to heal or prevent the formation ofnearly all cracks in a manner similar to theGTA welds in Heat 6. In Fig. 6, it is pri-marily the results from this heat that led usto incorporate the sulfur dependence atthe higher Creq/Nieq ratio of ~1.7. It is notclear why the solidification behavior ofHeat 9 is different than the other heatswith the same Creq/Nieq ratio, but, as withthe GTA welds, S appears to be playing arole as an apparent austenite former. Themixed mode solidification behavior of theother two heats with Creq/Nieq ratios of~1.74 is also different than the single-phase ferrite mode observed by Lienert,see Fig. 6. These differences further sug-

gest the possible role of sulfur in solidifi-cation behavior. However, it should benoted similar discrepancies in solidifica-tion mode at constant Creq/Nieq ratioswere also observed in another study in-volving austenitic stainless steels and elec-tron beam welds (Ref. 20), although in thecurrent study, the basic compositions arealmost identical except for sulfur. Differ-ences in weld parameters between the dif-ferent studies discussed here may also beplaying a role in solidification behavior(Ref. 52).

In light of the GTA Varestraint test re-sults, it is interesting to note the laserwelds with the highest Creq/Nieq ratiosseemed resistant to cracking. This lack ofcracking is inconsistent with the hypothe-sis that cracking susceptibility is high forsingle-phase ferrite solidification whenthe transformation to austenite occurs atlower temperatures, as it would for a mas-sive transformation. It was evident thateven though the solubility of sulfur ishigher in ferrite than in austenite, a largeamount of sulfides still formed during so-lidification. Since no cracking was ob-served at any level of sulfur tested, a sim-ple eutectic healing mechanism cannot beinvoked.

In one of the heats with the lowestCreq/Nieq ratio that solidified as austeniteand exhibited extensive cracking (Heat 7),it was found low melting phases associatedwith the cracks contained both sulfur andphosphorus — Fig. 9. The sulfur contentof this heat was 0.11%. However, sulfideswere not observed in TEM in the austen-ite solidified welds of Heat 1 with thelower sulfur level, which still exhibited se-vere cracking. Thus, it is conceivable phos-phorus is primarily responsible for crack-

Fig. 18 — Summary of GTA weld results showing region of good solidi-fication cracking resistance.

Fig. 19 — Combined results of both GTA and pulsed laser welds (Figs.6 and 18) showing a region of good weldability limited by LBW at lowCreq/Nieq ratios and by GTA at high ratios. Note that only LBW data isshown.

P+S

(w

t-%

)

P+S

(w

t-%

)

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ing and not sulfur. The phosphorus con-tent of all the heats was relatively high,0.03%. Given the higher phosphorus solu-bility in ferrite than austenite, higher lev-els of phosphorus may be tolerated inthose heats solidifying as ferrite. In addi-tion, the finer grain size of primary ferritesolidification (Fig. 17) should also reducethe tendency for cracking.

The results of the GTA and laser weld-ing trials can be combined to describealloy compositions that are amenable towelding with either conventional or HEDprocesses. This combination is shown inthe summary plot of Fig. 19. As shown, alower weldability limit occurs at aCreq/Nieq ratio of ~1.7 and is establishedby the laser weld process, while the upperweldability limit of ~ 1.9 is established bythe GTA process. The boundaries at theselimits are sloped, which implies the limitsare somewhat dependent upon sulfur con-tent, although additional data is desirable.As discussed earlier, it is important to em-phasize that although the data is portrayedin terms of total impurity level (P+S), itshould be remembered this portrayal isbased more on convention than a strongexperimental or theoretical basis. Never-theless, Fig. 19 illustrates reasonable fab-rication weldability can probably beachieved in free-machining gradesthrough suitable choice of alloy composi-tion. The mechanical properties of GTA,EB, and laser beam welds in these alloysare reported elsewhere (Refs. 25, 26).

Summary

The cracking response of GTA weldswas controlled primarily by solidificationmode. Using Suutula equivalents, thelower bound of Creq/Nieq ratio for goodcracking resistance was ~1.6, while theupper bound was ~1.9. This response isconsistent with existing rationale for de-scribing solidification cracking behaviorof austenitic stainless steels. However, atthe upper limit, high sulfur appeared to re-duce the cracking susceptibility of the sin-gle-phase ferrite solidified welds. This re-duction may be attributed to eutectichealing involving sulfur-containing-liquidand/or a change in solidification mode dueto sulfur acting as a Nieq. The exact role ofsulfur appears to be complex. It was some-what surprising the heat with the highestCreq/Nieq ratio and the lowest level of sul-fur was the most susceptible to GTA weldcracking.

The solidification behavior of pulsedlaser welds was also generally related toCreq/Nieq ratio. However, solidification inthe laser welds is complex, and in manycases the welds exhibited mixed mode so-lidification behavior. As with the GTAwelds, the solidification cracking behavior

of the laser welds was related to solidifica-tion mode. At Creq/Nieq ratios of 1.55,welds solidified as austenite and were sus-ceptible to cracking. Welds in alloys withCreq/Nieq ratios of 1.74 solidified in amixed mode, with some regions solidifyingas ferrite and other regions as austenite,but, in general, appeared to exhibit goodcracking resistance. However, with thesame Creq/Nieq ratio but a higher level ofsulfur, 0.27%, the amount of austenite so-lidification increased and a small amountof cracking was observed. At the Creq/Nieqratios of 1.92, the welds solidified com-pletely as ferrite, but, unlike the GTAwelds, the resistance to cracking was highat all impurity contents. In the laser welds,phosphorus is likely to play a major role inweld cracking behavior. For the alloysstudied, good cracking resistance was ob-served within a range of Creq/Nieq ratios of~1.7 to 1.9, with the lower limit estab-lished by the laser process and the upperlimit established by the GTA cracking be-havior.

Acknowledgments

Special thanks are given to AndyGardea for metallographic services and toAnnette Newman for laboratory supporton weldability testing. Thanks are alsogiven to Jerry Knorovsky for initial assess-ment of the laser weld performance and toMike Cieslak for reviewing the manu-script. The experimental alloys were pro-duced by Carpenter Technology Corp.,and its participation is greatly appreci-ated. Sandia is a multiprogram laboratoryoperated by Sandia Corp., a LockheedMartin company, for the United StatesDepartment on Energy under ContractDE-AC04-94AL85000.

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