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Introduction

Solidification Cracking andWeldability Diagrams

Solidification cracking is a well-known weldability issue withaustenitic stainless steels. Solidifica-tion cracking is complex with bothmetallurgical and mechanical aspects(Ref. 1). In austenitic stainless steels,impurities of sulfur and phosphorusform low melting temperature eutec-tic-like films that can wet solidifica-tion grain boundaries, creating acrack-susceptible microstructure (themetallurgical aspect) at the terminalstages of solidification. When a crack-

susceptible microstructure is present,if the strain from the welding thermalcycle exceeds the strain tolerance ofthe microstructure (the mechanical as-pect), cracks can propagate throughthe last-to-solidify liquid. In austeniticstainless steel weld metal, the metal-lurgical aspects of solidification crack-ing have been shown to be a functionof both alloy composition and impuri-ty content. It was first observed that welds inaustenitic stainless steels containingsmall amounts (greater than 5%) offerrite at room temperature had high-er resistance to solidification cracking(Refs. 2, 3). Thus, the ability to pre-dict room temperature microstruc-ture of stainless steel weld metals has

been studied considerably in the past,with constitution diagrams developedby several groups of researchers(Refs. 4–10). Constitution diagramswere developed to predict weld metalferrite content based on chemicalcomposition to mitigate solidificationcracking. The constitution diagramsgroup elements that stabilize theaustenite phase into a nickel equiva-lent (Nieq) and those that stabilize theferrite phase into a chromium equiva-lent (Creq). After the establishment ofconstitution diagrams, Suutala andco-workers then related the Creq/Nieq

ratio to solidification mode (Refs. 11,12) for these types of alloys. The alloycomposition, defined by Creq/Nieq,was found to determine the primarysolidification mode, either primaryferrite or primary austenite. At lowCreq/Nieq values, the liquid transformsto austenite during solidification,hence primary austenite solidificationmode. As Creq/Nieq increases abovesome critical value, primary ferritesolidification occurs, where the firstsolid to form is ferrite that then un-dergoes solid-state transformation toaustenite. Ferrite remains at roomtemperature due to the rapid coolingrate during welding, preventing com-plete transformation to austenite. Furthering the development thatsolidification cracking resistance wasrelated to the solidification mode, Ku-janpää and coworkers related solidifi-cation cracking to both solidificationmode and impurity content (Ref. 13).The weldability diagram they devel-

Laser Weldability of 21Cr­6Ni­9Mn Stainless Steel:Part I – Impurity Effects and Solidification Mode

The relationship between solidification cracking susceptibility and chemical composition is examined for laser­welded Type 21Cr­6Ni­9Mn stainless steels

BY S. B. TATE, D. A. JAVERNICK, T. J. LIENERT, AND S. LIU

ABSTRACT For laser welded Type 21Cr­6Ni­9Mn (21­6­9) stainless steels, the relationship betweensolidification cracking susceptibility and chemical composition was examined, and primarysolidification mode (PSM) diagrams were developed to predict the solidification mode.Sigmajig testing was used with experimental heats of Type 21­6­9 to determine the effectof P and S on solidification cracking when primary austenite solidification occurred. Phos­phorus showed a larger influence on solidification cracking relative to S, and a relationshipof (P + 0.2S) was found for total impurity content. The PSM diagrams to predict solidifica­tion mode were developed by analyzing welds made at three travel speeds for a widerange of 21­6­9 alloys and some other similar alloys. The minimum Creq/Nieq required forprimary ferrite solidification increased as travel speed increased, with more alloys show­ing primary austenite solidification at higher travel rates. As travel speed increased from21 to 85 mm/s, the average solidification rate increased from 6 to 25 mm/s.

KEYWORDS • Laser Welding • Austenitic Stainless Steel Solidfication Mode • Solidifcation Cracking

S. B. TATE (stephen.tate@aksteel.com), formerly a graduate student at Colorado State, is with AK Steel Corp., Middletown, Ohio. D. A. JAVERNICK(daj@lanl.gov), and T. J. LIENERT (lienert@lanl.gov) are with Los Alamos National Laboratory, Los Alamos, N.Mex. S. LIU (sliu@mines.edu) is withColorado School of Mines, Golden, Colo.

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oped for arc welds mapped crackingbehavior on a plot of impurity contentof P plus S vs. Creq/Nieq. In the weld-ability diagram at low Creq/Nieq, belowapproximately 1.5, primary austenitesolidification occurs, and cracking waspresent unless the combined impuritycontent was below approximately 0.02wt-%. Above 1.5 Creq/Nieq, where pri-mary ferrite solidification occurs,cracking is prevented regardless of im-purity content. The weldability dia-gram was developed using arc welded300 series stainless steels. The de-crease in solidification cracking sus-ceptibility when primary ferrite solidi-fication occurs was later revealed to bea function of differences in the natureof the solidification boundary, wettingpropensity of the boundaries, and sol-ubility of impurities (Refs. 14–16). Laser welding offers advantages ofboth reduced heat input compared toarc welding, and high travel speeds in-creased production rates for manufac-turing austenitic stainless steel com-ponents. As high-energy-density weld-ing processes such as laser welding andelectron beam welding became morewidely used, it was soon discoveredthat the rapid solidification conditionsof high-energy-density welding at highsolidification rates in austenitic stain-less steels can cause a shift to primaryaustenite solidification when primaryferrite solidification mode would beexpected (Refs. 17–22). The high so-lidification rates increase undercoolingat the solidification front, increasingthe stability of the austenite relativeto the ferrite (Refs. 23, 24). The shift

in solidification behavior also changesthe composition ranges that would becrack susceptible (Refs. 23, 25), whichled to the development of weldabilitydiagrams for pulsed laser welding of300 series austenitic stainless steels(Refs. 23, 26, 27). The Lienert diagram(Ref. 26) showed the critical Creq/Nieq

for primary ferrite solidification shift-ed from 1.5 for arc welding, to approx-imately 1.7 for pulsed laser welding.To avoid cracking with primaryaustenite solidification, combined P + S impurity content less than ap-proximately 0.02 wt-% was required,similar to the arc welding diagram. Using alloy composition to shift thesolidification mode to primary ferriteis the most common strategy to miti-gate solidification cracking issues.However, impurity content control canalso be considered. In some applica-tions primary austenite solidificationmay be desired, such as when ferrite isdetrimental to mechanical or magneticproperties, or primary austenite maybe unavoidable due to material orprocess parameter constraints. Whenprimary austenite solidification can-

not be avoided, the impurity levels ofthe starting material must be man-aged to mitigate hot cracking. Controlof S and P levels occurs during steelproduction and is beyond the controlof the end user. However, knowing theeffects of P and S on solidificationcracking can provide guidance in ac-cepting or selecting materials based oncomposition. Several researchers haveshown that P has a larger effect on in-creasing solidification cracking suscep-tibility than S in both austenitic stain-less steels (Refs. 28–31) and high-manganese steels (Ref. 32). For a con-stant total S plus P content of 0.032wt-% with primary austenite solidifi-cation, Li and Messler (Ref. 31) ob-served larger crack lengths with vare-straint testing for high P heats relativeto high S heats. It is thought the Ptends to segregate more strongly, lead-ing to liquid enriched in P, comparedto S forming sulfides, which removethe sulfur from the liquid.

High­Nitrogen, High­Manganese Stainless Steels

Fig. 1 — Sigmajig test fixture in the laser weldingenclosure.

Fig. 2 — Schematic Sigmajig sample with weld overfill and bead morphology shown:A — after welding; B — after sectioning showing the three transverse cross sectionstaken for microstructural analysis. All dimensions in mm.

A B

Fig. 3 — As­tested Sigmajig sample welded at 21 mm/s travel speed: A — Photograph;B — Stereoscope image of cracking at the end of weld. The box in 3A indicates the areafor the image of 3B.

A B

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High-nitrogen, high-manganeseaustenitic stainless steels are a subsetof austenitic stainless steels that wereinitially developed out of the concernfor nickel shortages (Ref. 33). Nitro-gen is a strong austenite stabilizer thatallows reduced nickel content as wellas providing interstitial solid-solutionstrengthening. High manganese con-tents promote nitrogen solubility, re-ducing nitride precipitation, which en-hances both mechanical propertiesand corrosion performance (Ref. 33).Type 21Cr-6Ni-9Mn (21-6-9) high-ni-trogen, high-manganese austeniticstainless steel (also known as Nitron-ic® 40) is the focus of this work. Type21-6-9 steel offers improved mechani-cal properties relative to 300 seriesstainless steels with no loss of corro-sion resistance (Ref. 8). As with otheraustenitic stainless steels, solidifica-tion cracking can cause weldability is-sues with 21-6-9. Brooks (Ref. 34)showed that 21-6-9 stainless steel be-haves similar to 300 series stainlesssteels with regard to solidificationcrack susceptibility decreasing withprimary ferrite solidification, regard-less of impurity content for arc weld-ing. A previous study on laser weldinga limited range of chemical composi-tions of 21-6-9 has shown that a shiftin solidification mode under the rapidsolidification conditions of laser weld-ing does occur (Ref. 35). Thus the abil-ity to predict solidification mode as afunction of chemical composition forlaser welding would be beneficial. It isunlikely that the diagrams developedfor 300 series stainless steels are perti-nent to solidification cracking of 21-6-9. Given the lower N and Mn contentof 300 series stainless, the relation-ship between solidification crackingand chemical composition of 21-6-9 islikely different than predicted by exist-ing weldability diagrams. Brooks (Ref.34) also revealed that the low-melting-temperature films responsible for thecracking in 21-6-9 are enriched in P.The P-enriched films on the crack sur-face may indicate that P is more detri-mental to cracking than S in 21-6-9. The relationships between alloycomposition, impurity content, solidi-fication mode, and solidification crack-ing behavior of 21-6-9 under high-en-ergy-density welding are not well char-acterized. The goal of this work was to

characterize the solidification modeand solidification cracking behavior of21-6-9 during laser welding for a widerange of chemical compositions. InPart I of this investigation, the effectsof S and P on solidification crackingwith primary austenite was studied,and PSM diagrams relating solidifica-tion mode to chemical compositionwere developed using a large numberof heats of 21-6-9. The PSM diagramswere developed for a range of solidifi-cation rates to observe the shift incritical chemical composition for pri-mary ferrite solidification as solidifica-tion rate increases. Part II will presentweldability diagrams that were devel-oped relating solidification mode andsolidification cracking to chemicalcomposition and impurity content.

Experimental

Materials

A variety of alloys and heats of high-nitrogen, high-manganese stainlesssteel materials were collected. The ma-terials used were of three different cate-gories: commercial wrought heats of 21-6-9, commercial wrought heats ofother high-nitrogen, high-manganeseaustenitic stainless steels, and experi-mental heats of 21-6-9 alloys produced

by laboratory melting. The chemicalcompositions of all the materials usedare given in Table 1. The focus of thiswork was 21-6-9 alloys, but it was desir-able to compare the applicability of thediagrams developed for a few otherhigh-nitrogen, high-manganese alloys.All chemical compositions were deter-mined using optical emission spec-troscopy (OES) for the majority of theelements, and Leco inert gas fusiontechniques for nitrogen, carbon, andsulfur. The OES instrument was cali-brated using a Nitronic 40 standard,IARM 19B. The results shown are theaverage of three analyses on each heatfrom the same laboratory. Experimental 21-6-9 heats weremelted for two purposes. One was to ex-amine the individual effects of sulfurand phosphorus on solidification crack-ing while isolating solidification modeto primary austenite (Alloys 40–53),and the other was to increase the com-position range of Type 21-6-9 beyondthe commercial heats but still withinthe nominal composition range of 21-6-9 (Alloys 54–59). Alloys 40–53 withvarying levels of sulfur and phosphoruswere melted from one starting alloycomposition. One commercial 21-6-9heat with low impurity levels that alarge enough quantity was available (Al-loy 14) was used as the starting masteralloy for the impurity variation heats.The experimental heats were inductionmelted under controlled atmospherewith a heat size of approximately 130 g.Additions of elemental sulfur and phos-phorus were used to modify the impuri-ty content. Solidification mode was con-trolled through additions of Ni andvarying the nitrogen pressure on themelt. The 10-mm-thick ingots were hot

Fig. 4 — Cross section of Sigmajig sam­ples showing the extent of solidificationcracking for A — Alloy 50; B — Alloy 53.Note the difference in magnification andthe arrow indicating cracking in Alloy 53.

A B

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rolled from 1200°C in three passes to afinal thickness of 5 mm. After two addi-tional cold rolling passes to a final thick-ness of 4 mm, a final solutionizing an-neal at 1100°C for 30 min was followedby a water quench.

Laser Welding

Sigmajig weldability testing (Ref.36) was used to compare the solidifica-tion crack susceptibility of the variousalloys. The Sigmajig test applies a pre-loaded tensile stress to a sheet samplewhile the sample is welded autoge-nously transverse to the tensile stressdirection. Figure 1 shows the Sigmajig fixturewith sample after welding. The stress

is applied by compressing two stacksof Bellville washers attached to themoving jaw. The load is controlled bymonitoring the stress in the instru-mented bolts that compress the wash-er stack. Quantifying the solidificationcracking for a given stress allows a rel-ative comparison of weldability forvariations in chemical composition orprocessing parameters. Samples 32 24 2 mm were used with the stressapplied along the 32 mm length andwelding along the 23 mm length. Astress of 310 MPa (45 ksi) was used forall samples. This stress level was cho-sen based on preliminary testing on aknown crack-susceptible alloy (Type309 stainless steel with primaryaustenite solidification and 0.037

wt-% P + S) as the minimum stresslevel to consistently cause cracking un-der the chosen processing parameters. The welding power source was a 1-kW multimode IPG fiber laser with100-m process fiber, 120-mm colli-mator, and 200-mm focal length lens,giving a theoretical spot size of 167m. All welding was done with opticalfocus (determined by Kapton filmburns) on the material surface and Arshielding gas flowing at 12 L/minthrough a nozzle parallel to the weld-ing direction with the nozzle leadingthe weld pool. The laser optics weretilted 7 deg relative to the travel direc-tion with the laser leading to reducethe chance of back reflections. Threetravel speeds of 21, 42, and 85 mm/s

Table 1 — Alloy Compositions (wt­%, balance Fe)

Alloy Type Element Cr Ni Mn N C Si P S Mo Ti Nb Cu V Al Co

1 21­6­9 19.90 6.49 8.94 0.23 0.031 0.54 0.017 0.001 0.047 0.001 0.019 0.010 0.039 0.025 0.193

2 21­6­9 19.89 6.55 8.96 0.31 0.030 0.49 0.016 0.001 0.100 0.001 0.019 0.016 0.039 0.032 0.205

3 21­6­9 19.80 7.11 9.43 0.27 0.030 0.39 0.019 0.001 0.179 0.001 0.026 0.000 0.099 0.023 0.055

4 21­6­9 19.93 6.51 8.93 0.30 0.032 0.50 0.017 0.001 0.103 0.001 0.019 0.020 0.040 0.034 0.195

5 21­6­9 19.88 6.14 8.91 0.28 0.029 0.49 0.018 0.001 0.102 0.001 0.022 0.052 0.136 0.029 0.158

6 21­6­9 19.46 7.26 9.32 0.27 0.021 0.54 0.019 0.001 0.111 0.001 0.022 0.182 0.053 0.021 0.149

7 21­6­9 19.40 7.20 9.04 0.26 0.034 0.45 0.020 0.001 0.092 0.001 0.032 0.137 0.137 0.022 0.096

8 21­6­9 19.74 6.62 9.47 0.27 0.021 0.43 0.027 0.005 0.382 0.008 0.056 0.343 0.132 0.014 0.317

9 21­6­9 19.73 6.44 8.72 0.31 0.021 0.31 0.022 0.001 0.080 0.001 0.023 0.159 0.050 0.010 0.040

10 21­6­9 20.44 7.06 8.94 0.36 0.028 0.70 0.023 0.001 0.248 0.001 0.041 0.350 0.063 0.025 0.081

11 21­6­9 19.43 6.20 9.36 0.33 0.034 0.38 0.023 0.002 0.227 0.001 0.034 0.243 0.112 0.009 0.066

12 21­6­9 18.74 6.89 9.31 0.23 0.011 0.30 0.019 0.002 0.044 0.001 0.019 0.000 0.027 0.024 0.040

13 21­6­9 19.49 6.89 9.31 0.30 0.028 0.36 0.023 0.001 0.419 0.001 0.038 0.197 0.096 0.020 0.161

14 21­6­9 19.99 6.35 8.88 0.27 0.032 0.44 0.021 0.002 0.126 0.006 0.055 0.121 0.130 0.007 0.136

30 18­2­12 18.64 1.23 11.50 0.35 0.101 0.71 0.021 0.001 0.088 0.001 0.031 0.105 0.094 0.009 0.049

31 SCF­260 18.73 2.73 17.40 0.60 0.035 0.54 0.024 0.001 2.258 0.001 0.028 0.040 0.149 0.027 0.063

32 15­15HS Max 18.89 1.02 16.97 0.57 0.033 0.43 0.027 0.001 0.816 0.001 0.033 0.085 0.161 0.023 0.070

33 Nitronic 50 21.42 15.18 4.97 0.35 0.048 0.35 0.020 0.007 2.466 0.056 0.127 0.093 0.199 0.011 0.135

34 Nitronic 50 21.23 11.92 5.00 0.25 0.031 0.27 0.024 0.004 2.090 0.005 0.140 0.500 0.140 0.005 0.050

35 Nitronic 60 16.99 7.95 7.78 0.16 0.061 3.74 0.032 0.001 0.341 0.010 0.005 0.273 0.065 0.020 0.097

36 Nitronic 30 16.58 3.28 8.51 0.17 0.029 0.33 0.024 0.005 0.110 0.005 0.013 0.523 0.059 0.008 0.071

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(50, 100, and 200 in./min) were used.The laser power was adjusted at eachtravel speed to maintain a complete-joint-penetration weld on the 2-mmsample thickness. The laser power,measured with an Ophir Comet 1K,was 555, 755, and 1166 W at 21, 42,and 85 mm/s travel speed, respective-ly. Samples were ultrasonically cleanedin methanol prior to welding to re-move surface contaminants.

Microstructural Analysis

The surfaces of the Sigmajig sam-

ples were inspected for cracking at upto 100 magnification after welding.After examining the surface, threetransverse cross sections were re-moved from the Sigmajig samples atapproximately 3 and 6 mm from theend of the weld as shown in Fig. 2. Thetransverse cross sections were all tak-en at the end of the weld because thatis the location where cracks were ob-served in the surface inspection whenpresent. Thermomechanical modelingresults of Sigmajig testing from litera-ture supports the choice of locationfor cross sections, showing that thetransverse tensile stresses are all de-

veloped toward the end of the weld(Ref. 37). All three cross sections were exam-ined to characterize solidificationmode and solidification crackingthrough microstructural observation.Standard metallographic preparationprocedures of grinding with SiC andpolishing with diamond compoundswere used, and electrolytic etching wasdone with 10% oxalic acid. The solidi-fication mode was characterized withlight optical microscopy at a variety ofmagnifications. Examination for solid-ification cracks was also conductedwith light optical microscopy with a

Table 1 — Continued

Alloy Type(a) Element

Cr Ni Mn N C Si P S Mo Ti Nb Cu V Al Co

40 14 19.67 7.06 8.86 0.39 0.030 0.46 0.022 0.003 0.127 0.006 0.055 0.136 0.138 0.011 0.138

41 14 19.65 7.14 8.95 0.33 0.030 0.46 0.023 0.012 0.127 0.006 0.055 0.138 0.137 0.011 0.142

42 14 19.67 6.98 8.80 0.34 0.031 0.47 0.021 0.024 0.125 0.006 0.055 0.140 0.141 0.011 0.140

43 14 19.46 6.97 9.06 0.33 0.031 0.48 0.023 0.035 0.127 0.006 0.055 0.149 0.144 0.011 0.137

44 14 19.60 7.02 8.74 0.38 0.030 0.45 0.025 0.003 0.125 0.006 0.055 0.131 0.133 0.011 0.132

45 14 19.69 7.11 8.74 0.37 0.031 0.46 0.030 0.003 0.127 0.006 0.055 0.136 0.139 0.011 0.134

46 14 19.58 7.09 8.91 0.38 0.030 0.47 0.035 0.003 0.129 0.006 0.055 0.139 0.140 0.011 0.134

47 14 19.53 7.13 9.00 0.38 0.031 0.46 0.031 0.012 0.128 0.006 0.055 0.137 0.137 0.011 0.130

48 14 19.54 7.08 8.81 0.37 0.030 0.44 0.024 0.009 0.125 0.006 0.055 0.128 0.131 0.011 0.134

49 14 19.39 7.06 9.04 0.33 0.031 0.47 0.026 0.019 0.128 0.006 0.055 0.142 0.140 0.011 0.135

50 14 19.52 7.05 8.87 0.37 0.030 0.48 0.031 0.019 0.128 0.006 0.055 0.145 0.143 0.011 0.137

51 14 19.33 6.90 8.98 0.35 0.030 0.48 0.024 0.041 0.125 0.006 0.055 0.153 0.144 0.011 0.140

52 14 19.41 7.02 8.93 0.31 0.031 0.45 0.022 0.048 0.128 0.006 0.055 0.135 0.136 0.011 0.135

53 14 19.35 6.98 8.97 0.32 0.029 0.47 0.023 0.058 0.128 0.006 0.055 0.142 0.140 0.011 0.138

54 11 18.93 7.33 9.64 0.36 0.034 0.42 0.034 0.012 0.215 0.004 0.031 0.258 0.114 0.014 0.069

55 11 18.97 7.35 9.40 0.36 0.034 0.42 0.034 0.017 0.214 0.003 0.030 0.249 0.113 0.012 0.067

56 14 19.88 6.16 9.19 0.34 0.030 0.46 0.029 0.014 0.125 0.002 0.018 0.140 0.131 0.016 0.136

57 14 19.84 6.24 9.14 0.34 0.031 0.48 0.035 0.019 0.127 0.002 0.019 0.147 0.135 0.016 0.138

58 14 20.01 6.16 9.21 0.31 0.031 0.48 0.031 0.012 0.125 0.002 0.019 0.147 0.135 0.017 0.135

59 14 20.12 6.18 8.93 0.27 0.031 0.45 0.033 0.019 0.124 0.001 0.017 0.136 0.130 0.014 0.132

(a) Type for Alloys 40–59 indicate the starting material used for melting.

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maximum of 1000 magnification. Forchemical composition analysis, EDSwith a JEOL JSM-7000F SEM wasused. Correlation of the solidificationmode, solidification cracking response,and the chemical composition wasused to develop the weldability dia-grams for the three travel speeds used.Quantifying the length of cracks(when present) was used to assess theseverity of solidification cracking foreach cross section.

Results and Discussion

Effect of Sulfur and Phosphoruson Solidification Cracking

The target impurity levels of S and Pfor Alloys 40–53 are given in Table 2.Sulfur and phosphorus were independ-ently increased to levels that causedcracking (the first row and first columnof Table 2), as well as combinations of Sand P, to create a three by three full fac-torial matrix. The full factorial sampleswere used to analyze the possibility ofinteractions between S and P. The crack-ing behavior, impurity content, and so-lidification mode of all alloys and alltravel speeds are given in Table 3. OnlyAlloys 40–53 welded at 21 mm/s travel

speed were used to analyze the effectsof P and S on solidification cracking. Re-sults are given for Alloys 40, 49, and 53at higher travel speeds because they areincluded in higher travel speed weldabil-ity diagrams. For Alloys 40–53, there issome unintentional variation in Creq/Nieq due to differences in Ni and N addi-tions, but the solidification mode wasisolated to primary austenite as desired. Figure 3 shows an as-tested Sigma-jig sample from 21 mm/s travel speed.Surface cracking, if present, was onlyobserved at the end of the weld. Sub-surface cracking was often presentwith no indication of cracking on theweld surface. Alloys 46, 47, 49, 50, 52,and 53 all exhibited cracking. Crackingin the alloys with no S added occurredat a level of 0.035 wt-% P, as in Alloy46. When only S was added, 0.048 wt-% S was required to promote crack-ing, and with a much higher residual Pcontent compared to the residual Slevel in the alloys with P added. Amuch higher P plus S level was re-quired to initiate solidification crack-ing for a high-S, low-P alloy as com-pared to low-S, high-P alloy. Amongthe alloys with cracking observed, alarge variation in crack length is pres-ent. Alloys 46 and 50, with the highestP contents, exhibited the largest crack

lengths, with large through-thicknesscenterline cracks. Figure 4 shows thedifference in crack length for high-Pcompared to high-S, from Alloys 50and 53 with a large centerline crackpresent in Alloy 50 and a small non-centerline crack on the right side ofthe weld for Alloy 53. Both the impurity level at crack ini-tiation and the extent of the cracklengths indicated that P has a largerinfluence on solidification crack sus-ceptibility than S. Regression analysiswas used to quantify the different ef-fects of P and S on cracking. Responsesof the total number of cracks for an al-loy, and number of sections withcracks for factors of P, S, and P times Swere used to run general factorial re-gression on the three by three matrix(Alloys 40–42, 44, 45, and 47–50). Thep-values for the various factors and re-sponses are given in Table 4. The p-values for phosphorus are both below0.05, indicating P has a statisticallysignificant effect for both responses.Sulfur showed slightly higher p-valuesthan phosphorus overall, and for theresponse of the number of crackswhere the p-value is greater than 0.05,the level of S is likely not significantfor that response. The p-values for theP times S interaction term showed thehighest values, and both are greaterthan 0.05, indicating there was no ef-fect of interaction between P and S onsolidification cracking. Knowing there were no P-S interac-tion effects, a general linear regressionmodel was generated using Alloys40–53 for responses of the number ofcracks per alloy and number of cross-sections showing cracking. Despite thehigh p-value for S in the model based onthe number of cracks, it was expected tobe significant without the interactionterm. With the interaction term re-moved, both P and S were significantfor both responses analyzed. The result-ing model for the number of cracks was

Number of Cracks = –3.90 + 151.7 P (wt-%) + 31.9 S (wt-%).

Similarly, for the number of crosssections showing cracks, the modelwas

Number of Sections with Cracks = –5.38 + 216.3 P (wt-%) + 32.2 S (wt-%).

The residual plots for both models

Fig. 5 — Chemical composition of the liquid as a function of fraction of solid for A — Sul­fur and B — Phosphorus using Thermo­Calc Scheil simulation for 0.02 wt­% P and 0.06wt­% S nominal composition under primary austenite solidification mode.

Table 2 — Levels of Sulfur and Phosphorus in Experimental Heats (wt­%)

Phosphorus Sulfur 0.003 0.012 0.024 0.035 0.041 0.048 0.058

0.022 40 41 42 43 51 52 53 0.025 44 48 49 — — — — 0.030 45 47 50 — — — — 0.035 46 — — — — — —

A B

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indicated the models were valid. Nor-malizing the coefficients for P and S ineach model to a coefficient of one for P,gives a coefficient for S of 0.21 basedon the number of cracks, and 0.15based on the number of sections withcracks. Both models indicate a reducedeffect of S relative to P, similar to re-sults from literature for effects of S andP on cracking in other austenitic stain-less steels (Refs. 28–31). Based on theregression analysis, a coefficient of 0.2for S was chosen for the weldability dia-

grams presented in Part II. To better understand the partition-ing behavior of P and S, and phaseslikely to form during solidification,Scheil solidification simulations forthe primary austenite solidificationmode were done using Thermo-Calcwith the TCFE7 database. The resultsare presented for three compositions:a low-S, low-P alloy, a low-S, high-P al-loy, and a high-S, low-P alloy. Sulfurand phosphorus levels were chosen tosimulate Alloys 40, 46, and 53. All

three calculations had constant levelsof all other elements used, with com-position adjusted to give primaryaustenite solidification. Evaluating thecomposition of the liquid at a highfraction of solid should give an idea ofthe enrichment of P and S in the finalto solidify liquid to which solidifica-tion cracking is attributed. The start-ing composition of P and S in the sim-ulated alloys and the composition ofthe liquid at 0.97 fraction of solid aregiven in Table 5.

Table 3 — Solidification Mode of Sigmajig Samples

Espy Equivalents Solidification Mode

Alloy P + 0.2 S Creq Nieq Creq/ Nieq 21 mm/s 42 mm/s 85 mm/s

1 21.03 12.33 1.71 0.018 F F F 2 21.02 13.53 1.55 0.016 F D A 3 21.14 13.44 1.57 0.019 F F A 4 21.10 13.51 1.56 0.017 F F A 5 21.49 12.55 1.71 0.018 F F F 6 20.72 13.36 1.55 0.020 F A A 7 20.93 13.42 1.56 0.020 F A A 8 21.50 12.68 1.69 0.028 F F F 9 20.57 13.32 1.54 0.022 F F D 10 22.14 15.20 1.46 0.023 F A A 11 20.83 13.90 1.50 0.023 F A A 12 19.45 12.06 1.61 0.019 F A A 13 21.01 13.82 1.52 0.023 F A A 14 21.47 12.76 1.68 0.021 F F F

30 20.31 11.29 1.80 0.021 F F A 31 22.65 15.81 1.43 0.024 F D D 32 21.23 13.39 1.59 0.027 F D A 33 25.51 23.53 1.08 0.021 A A A 34 24.51 18.28 1.34 0.025 A A A 35 23.33 14.20 1.64 0.032 D A A 36 17.18 8.95 1.92 0.025 F F F

40 21.23 15.71 1.35 0.023 A A A 41 21.21 14.68 1.44 0.025 A — — 42 21.27 14.72 1.44 0.026 A — — 43 21.09 14.62 1.44 0.030 A — — 44 21.12 15.45 1.37 0.025 A — — 45 21.27 15.41 1.38 0.031 A — — 46 21.17 15.71 1.35 0.035 A — — 47 21.10 15.65 1.35 0.033 A — — 48 21.04 15.39 1.37 0.025 A — — 49 20.98 14.57 1.44 0.030 A A A 50 21.14 15.33 1.38 0.035 A — — 51 20.96 14.73 1.42 0.032 A — — 52 20.96 14.16 1.48 0.032 A — — 53 20.94 14.32 1.46 0.035 A A A 54 20.40 15.57 1.31 0.037 A A A 55 20.42 15.58 1.31 0.038 A A A 56 21.40 13.86 1.54 0.032 F D A 57 21.42 13.96 1.53 0.039 F D A 58 21.59 13.37 1.61 0.033 F D A 59 21.62 12.60 1.72 0.037 F D D

Solidification modes: A austenite, D dual (austenite and ferrite), F ferrite.

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At minimum P and S levels, the liq-uid is enriched in P to approximately 1wt-% while the S level reaches approxi-mately 0.03 wt-%. When the starting Plevel is doubled to 0.04 wt-%, the P inthe liquid almost doubles as well. Themuch larger increase in S content from0.002 to 0.06 wt-% showed almost nochange in the composition of S in theliquid. The low amount of S in the lastliquid to solidify, even at high initialcompositions, is attributed to the for-mation of MnS during solidification.The compositions of P and S in the liq-uid phase are plotted as a function offraction of solid in Fig. 5 for the high-S, low-P case. Phosphorus in the liquidcontinues to increase as fraction of sol-id increases, as would be expected for a

solute with a parti-tioning coefficientless than one. Thesulfur content of theliquid increases fromthe nominal compo-sition initially, andthen at approximate-ly 0.4 fraction of sol-id begins to decrease.The shift to decreas-ing S is due to theformation of man-ganese-sulfide(MnS), which isshown by the changein line color from redto green. As MnScontinues to formfrom the melt, the Scontent of the liquid

is actually reduced below the nominalcomposition. Similar behavior is ob-served for the low-S level, with the liq-uid being enriched with S until MnSforms at approximately 0.9 fraction ofsolid.

Based on the Scheil calculations of Pand S compositions in the final liquid,it would be expected that the cracksurfaces would exhibit P-rich films,and the S would be dispersed as MnSinclusions within the weld microstruc-ture. A crack in a cross section of Alloy47 is shown in Fig. 6, along with theEDS spectrum for a spot scan of thecrack surface. The crack showed a film-like second phase along the crack. TheEDS spectra showed no indications ofS or P, and a small Si peak that is typi-cal of the matrix. Other EDS attempts

on similar crack features on multiplesamples were unsuccessful in identify-ing the constituents along the cracks.The large interaction volume of thebeam compared to the size of the fea-ture being characterized likely makethe signal from the matrix overpowerthe characteristic x-rays from the crackfilm. Exposing the crack surface to in-crease the volume of film material wasalso unsuccessful in identifying the el-ements present at the crack surface.

An example of the exposed cracksurface is shown in Fig. 7. The cracksurface showed the egg crate morphol-ogy typical of solidification cracking.The film-like nature of the secondphase present can be seen between the dendrite tips. Further characteri-zation of the composition of the cracksurfaces with electron-microprobeanalysis (EMPA) or Auger electronspectroscopy would likely provide suit-able results, but was not conducted inthis work. It is important to note thatBrooks (Ref. 34) identified P- and Mn-rich films along the solidification grainboundaries in GTA welded 21-6-9using EMPA. Similar chemical compo-sitions are expected for the secondphases present along the cracks ob-served here.

Chemical Composition and Solidification Cracking

The larger increase in solidificationcrack susceptibility for P relative to S forprimary austenite solidification ob-served here is similar to that observedby Arata et al. (Refs. 28, 29), Katayamaet al. (Ref. 30), and Li and Messler (Ref.

Fig. 6 — Unetched crack in the cross section of Alloy 47 welded at 21 mm/s: A — Note film of secondary phase along crack; B — EDS spec­trum of spot scan indicated in A by red dot.

Fig. 7 — The dendritic fracture surface of Alloy 55 welded at 85mm/s with the second phase lining areas of the crack surface.

A B

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31). The literature results for 300 seriesstainless steels with austenitic solidifi-cation all indicate P being far moredetrimental to solidification cracking.Katayama (Ref. 30) showed phosphidestend to form film-like features along so-lidification grain boundaries, while sul-fides form with a globular morphology.Brooks (Ref. 34) showed that a Mn- andP-rich film is found on the partiallymelted grain boundaries for GTA weld-ing of 21-6-9 while S is tied up in dis-crete MnS particles. Similar results forthe morphology of phosphides and sul-fides would be expected for the experi-mental 21-6-9 alloys used here. Theelectron microscopy work shown in lit-erature was able to identify the P and Scompositions likely due to the muchlarger size of the phosphides and sul-fides in arc welding. Based on the literature and Scheilsolidification calculations performed,

the sulfides are expected to be MnStype. With MnS forming during solidi-fication even at very low (0.003 wt-%)S levels, it is expected that any S ex-ceeding the solubility of the liquid isprecipitated out during solidification.The melting temperature of MnS ofapproximately 1300°C (Ref. 30) ishigher than the terminal solidificationtemperature of 1230°C predicted fromthe Thermo-Calc calculations for theexperimental 21-6-9 Alloys, indicatingthe sulfides form prior to final solidifi-cation and prevent S from forming low

melting temperature phases that couldpromote solidification cracking. Giventhat film-like low melting temperaturefeatures are considered detrimental tosolidification cracking, the globularmorphology of MnS would also reducethe influence of S on solidificationcracking susceptibility. A coefficient of 0.2 for S, as deter-mined in the analysis of the experimen-tal 21-6-9 heats, may be related to thehigh Mn content of 21-6-9. Honey-combe and Gooch (Ref. 44) showed Mnlevels of approximately 2–9 wt-% were

Table 4 — p­Values for Responses in Regression Model

Term Number of Cracks Number of Sections with Cracks

P 0.012 0.009 S 0.193 0.021 P x S 0.262 0.057

Fig. 8 — Micrographs showing observed solidification modes: A — Primary austenite solidification (A) in Alloy 13 welded at 85 mm/s; B —primary austenite solidification with interdendritic ferrite (AF) and cracking in Alloy 35 welded at 21 mm/s; C — primary ferrite solidifica­tion (F) in Alloy 3 welded at 21 mm/s; D — dual solidification modes (D) with dark etching primary austenite solidifying epitaxially fromthe weld interface transitioning to lighter etching primary ferrite toward the centerline in Alloy 9 welded at 85 mm/s.

A B

C D

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effective in reducing solidificationcracking in a fully austenitic Type 310weld metal with 0.01 wt-% S and 0.02wt-% P, but no mechanism for the re-duction in crack susceptibility was giv-en. For alloys with lower Mn contentthan Type 21-6-9, such as 300 seriesstainless steels, it is possible that lower

S levels would initiate so-lidification cracking, giv-ing a higher S coefficientfor those alloys than 0.2as determined for Type21-6-9. One final consid-eration is that given thelow levels of S in com-mercial alloys used here(typically less than 0.005wt-%), a small variationin the coefficient for Smakes little difference inoverall calculated impu-rity content for commer-cial alloys.

PrimarySolidification Mode(PSM) Diagrams

The PSM diagrams were produced byplotting Nieq as a function of Creq. Toplot the Cr and Ni equivalencies of thealloys studied here, the equivalents de-veloped by Espy (Ref. 8) were used,which are given by the following equa-tions:

Creq =%Cr + %Mo + 1.5 × %Si + 0.5 × %Nb + 5 × %V + 3 × %AlNieq = %Ni + 30 %C + 0.87 for Mn+ 0.33 %Cu + (%N – 0.045) NCoef.

where NCoef. = 30 when nitrogen is0.0–0.20 wt-% or NCoef. = 22 when ni-trogen is 0.21–0.25 wt-% or NCoef.= 20 when nitrogen is 0.26–0.35 wt-%. Thecoefficients developed by Espy areknown to be the most applicable Crand Ni equivalents for arc welding ofhigh-N, high-Mn austenitic stainlesssteels. Ritter and coworkers (Refs. 38,39) determined that Espy equivalentsare the most applicable for arc weldingof Nitronic 50 alloys for both ferritecontent and solidification mode pre-diction. Ritter and Savage (Ref. 39)also plotted arc welding solidificationmode results for a large number ofhigh-N, high-Mn stainless steels fromSuutala (Ref. 40) showing the Espyequivalents provided a good fit. Robi-no et al. (Ref. 41) showed that theEspy equivalents were preferred for arcwelding of Nitronic 60 and anothersimilar alloy. Other equivalents were also ex-plored in developing the diagrams;specifically, Hull (Ref. 6) and Hammarand Svensson (Ref. 7) were considered.Among the three equivalents consid-ered, the Espy equivalents providedthe clearest trends of solidificationmode as a function of Creq/Nieq. TheEspy coefficients were chosen basedon better discrimination of the solidi-fication modes observed in this work,and their known applicability to arcwelding of Nitronic alloys. For the

Fig. 9 — PSM diagram for 21 mm/s (50 in./min) with Nieq asa function of Creq. A is the primary austenite solidification,and F is the primary ferrite solidification. Other indicatesnon­21­6­9 type.

Fig. 11 — PSM diagram for 85 mm/s (200 in./min) with Nieqas a function of Creq. A is the primary austenite solidifica­tion, F is the primary ferrite solidification, D is the dual withmixed primary austenite and primary ferrite solidification.Other indicates non­21­6­9 type.

Fig. 10 — PSM diagram for 42 mm/s (100 in./min) with Nieq as afunction of Creq. A is the primary austenite solidification, F is the pri­mary ferrite solidification, D is the dual solidification mode. Otherindicates non­21­6­9 type.

Table 5 — Thermo­Calc Scheil Simulation Liquid Compositions

Starting Composition Composition of Liquid at 0.97 (wt­%) Fraction Solid (wt­%)

P S P S 0.02 0.002 0.98 0.028 0.04 0.002 1.74 0.024 0.02 0.06 0.92 0.030

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PSM diagrams, the nitrogen loss thatoccurs during high-energy-densitywelding (Refs. 34, 35) was consideredwhen calculating the Nieq. Based onpreviously reported results (Ref. 42),nitrogen loss of 10% from the initialbase metal composition was used tocalculate the Nieq. Previously developed PSM andweldability diagrams do not accountfor variations in solidification condi-tions within a given welding process,whether developed for laser or arcwelding. It is known that process pa-rameters different from those used todevelop a weldability diagram cancause the diagram to be not applicable(Ref. 43). It was desired to incorporatethe shift in solidification mode thatoccurs due to changes in solidificationrate into this work. PSM diagramswere developed at three distinct travelspeeds that span a range common forcontinuous-wave laser welding. Theresult is three separate PSM diagramsthat capture the shift in solidificationmode with solidification rate. The Creq, Nieq, impurity content, andsolidification mode of the Sigmajig sam-ples used to generate the diagrams weregiven above in Table 3. The primary so-lidification modes were identified bytheir distinctive microstructures. Typi-cal solidification microstructures ob-served are shown in Fig. 8. Primaryaustenite microstructures with andwithout ferrite were observed. The cel-lular primary austenite solidificationwith no ferrite shown in Fig. 8A was themajority of the primary austenite solidi-fication observed. Less common was thedendritic primary austenite containingsmall amounts of ferrite shown in Fig.8B. A typical primary ferrite solidifica-tion microstructure is shown in Fig. 8C.No cracks were observed with primaryferrite solidification for all travelspeeds. Dual-mode solidification mi-crostructures with both primaryaustenite and primary ferrite were ob-served, shown in Fig. 8D. Dual solidifi-cation mode was observed to exhibitboth ferrite and austenite growing to-gether from the weld interface, or onemode transitioning to another as loca-tion changed within the weld. In dualsolidification mode alloys that exhibitedcracking, the cracking always occurredin the primary austenite regions. Pri-mary ferrite solidification followed bywhat has been referred to in literature

as massive decomposition of the ferriteto austenite (Refs. 15, 20, 23) that canoccur under rapid cooling was observedin some alloys at 85 mm/s travel speed,shown in Fig. 8D. The PSM diagram developed for 21mm/s (50 in./min) travel speed in-cludes a larger number of alloys due tothe large number of compositionsused for the impurity content varia-tion work. Only Alloys 40, 49, and 53were repeated at 42 and 85 mm/s trav-el speed. The PSM diagram for weldsat 21 mm/s (50 in./min) is shown inFig. 9 with Nieq as a function of Creq. AsNieq increases for a given Creq, the so-lidification mode shifts from primaryferrite to primary austenite solidifica-tion. For the 21-6-9 alloys tested,there appears to be a clear shift in so-lidification mode, indicated by thedashed line. The PSM diagram developed for 42mm/s (100 in./m) travel speed isshown in Fig. 10. As expected, a largernumber of alloys showed primaryaustenite solidification at the in-creased travel speed. A number of alloys also showed thedual solidification mode, which wasnot observed at 21 mm/s. The demar-cation between solidification modeswas less prominent. Note that all al-loys are shown on the plot, but the so-lidification mode boundary was drawnbased only on the 21-6-9 alloys(shown in black), because that was thefocus of this work. Figure 11 shows the PSM diagramdeveloped for 85 mm/s. Again, the so-lidification mode boundary shifteddownward as the travel speed increasedwith a larger number of alloys showingprimary austenite solidification. Similarto the 42 mm/s diagram, some of theother alloys show primary austenite so-lidification below the line indicating thechange to primary ferrite solidificationfor the 21-6-9 alloys.

Solidification Mode Considerations

The PSM diagrams can be comparedto the diagrams generated by Ritter andSavage (Ref. 39) with Espy coefficientsfor arc welding of similar alloys. Thetransition Creq/Nieq between solidifica-tion modes has shifted downward in thelaser welding diagrams presented herecompared to the arc welding diagram.

The downward shift in the solidificationmode boundary indicates the Creq/Nieq

required for primary ferrite solidifica-tion increased for the laser welding, asexpected. Within the laser welding PSMdiagrams, the solidification modeboundary shifted downward, thus theminimum Creq/Nieq for primary ferritesolidification increased as travel speedincreases. The mix of solidificationmodes that was found to occur close tothe boundary between solidificationmodes at higher travel speeds in thiswork is similar to the weldability dia-gram for pulsed laser welding of 300 se-ries alloys (Ref. 26), where both primaryferrite and dual solidification modeswere observed for Creq/Nieq values be-tween 1.6 and 1.7. There is no reason to assume thatthe boundary between solidificationmodes should be a line. For example,the liquidus lines of the Fe-Cr-Ni andFe-Cr-Mn systems show a curvedboundary between primary solidifica-tion modes over a large range of chemi-cal compositions. Based on the fit of theobserved solidification modes to the im-posed boundaries in this work andknowing that linear boundaries be-tween the solidification modes are anassumption, better equivalencies mayneed to be developed. The linear Cr andNi equivalencies also do not account forthe interaction terms that are presentwhen considering the thermodynamicsof the systems. Work to develop new Crand Ni equivalencies should considerboth of these aspects. Quantifying the change in the solidi-fication mode boundary line as travelspeed increased was attempted, but itwas considered inappropriate given thelack of clear fit of the lines to the ob-served solidification modes at 42 and85 mm/s travel speeds. Developing bet-ter equivalencies may allow quantifyingthe downward shift in solidificationmode boundary that occurs as travelspeed increases. The transition from primary ferriteto primary austenite solidification for agiven Creq/Nieq with increasing travelspeed has been ascribed to increasedundercooling at high solidification rates.Estimating the solidification rate (veloc-ity of the solid/liquid interface duringsolidification) of the welds for each trav-el speed was done using

R = Vcos()

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where V is the welding travel speed and is the angle between the dendritegrowth and weld travel direction (Ref.45). Average solidification rates of 6mm/s at 21 mm/s travel speed, 13mm/s at 42 mm/s travel speed, and 25mm/s at 85 mm/s travel speed were ob-served at 50% penetration depth,where the solidification rate is expectedto be the highest (Ref. 46). The solidifi-cation rate appears to be approximately30% of the travel speed for all the con-ditions, which agrees with average so-lidification rates reported for continu-ous-wave laser welding of super-austenitic stainless steel (Ref. 47).Pulsed laser weld solidification rates arereported to range from 42 to 142 mm/s(Ref. 48), thus the highest travel speedsused in this work had solidificationrates that approach those in pulsedlaser welding. Again, the solidificationrates given are averages; both solidifica-tion rate and temperature gradient varyspatially within the weld pool. The vari-ation of solidification rate and tempera-ture gradient throughout the weld like-ly causes the dual solidification modewhere primary austenite transitions toprimary ferrite or vice versa.

Conclusions

For 21-6-9 alloys with primaryaustenite solidification mode, P has alarger effect on increasing solidificationcrack susceptibility relative to S. Interac-tion between P and S was found to havean insignificant effect on cracking. A co-efficient of 0.2 was determined for S,with total impurity content calculatedas P + 0.2S. The S is thought to not con-tribute significantly to solidificationcracking because of the formation ofglobular sulfides during solidification. A range of 21-6-9 alloys were used todevelop PSM diagrams to predict solidi-fication mode for laser welding of 21-6-9. The PSM diagrams were developed at21, 42, and 85 mm/s travel speed. Astravel speed increased the minimumCreq/Nieq for primary ferrite solidifica-tion increased, which was attributed tothe increase in solidification rate astravel speed increased.

The authors would like to thankLos Alamos National Laboratory for fi-

nancial support of this graduate re-search work. The authors also thankDr. Graham McIntosh formerly of Car-penter Technology Corp., Dr. LuisGarza of AK Steel Corp., and Dr. JohnElmer of Lawrence Livermore NationalLaboratory for donating materials forthis work. The authors also acknowl-edge the NSF Center for IntegrativeMaterials Joining Sciences for EnergyApplications for the collaborative re-search opportunity.

1. Cross, C. 2005. On the origin of weldsolidification cracking. Hot Crack. Phenom.Welds. 2. Hull, F. C. 1967. Effect of delta ferriteon the hot cracking of stainless steel. Weld-ing Journal 46(9): 399-s to 409-s. 3. Brooks, J., and Lambert, F. J. 1978.The effects of phosphorus, sulfur and fer-rite content on weld cracking of type 309stainless steel. Welding Journal 57(5): 139-sto 143-s. 4. Schaeffler, A. L. 1949. Constitutiondiagram for stainless steel weld metal. Met.Prog. 56(11): 680–680b. 5. DeLong, W. T., Ostrom, G. A., andSzumachowski, E. R. 1956. Measurementand calculation of ferrite in stainless steelweld metal. Welding Journal 35(11): 521-sto 528-s. 6. Hull, F. C. 1973. Delta ferrite andmartensite formation in stainless steels.Welding Journal 52(5): 193-s to 203-s. 7. Hammar, O., and Svensson, U. 1977.Influence of steel composition on segrega-tion and microstructure during solidifica-tion of austenitic stainless steels. Proceed-ings: Solidification and Casting of Metals. pp.401–410. 8. Espy, R. H. 1982. Weldability of ni-trogen-strengthened stainless steels. Weld-ing Journal 61(5): 149-s to 156-s. 9. Siewert, T. A., McCowan, C. N., andOlson, D. L. 1988. Ferrite number predic-tion to 100 FN in stainless steel weld met-al. Welding Journal 67(12): 289-s to 298-s. 10. Kotecki, D. J., and Siewert, T. A.1992. WRC-1992 constitution diagram forstainless steel weld metals: A modificationof the WRC-1988 diagram. Welding Journal71(5): 171-s to 178-s. 11. Suutala, N., Takalo, T., and Moisio,T. 1979. The relationship between solidifi-cation and microstructure in austeniticand austenitic-ferritic stainless steel welds.Metall. Trans. A 10(4): 5–7. 12. Takalo, T., Suutala, N., and Moisio,T. 1979. Austenitic solidification mode inaustenitic stainless steel welds. Metall.

Trans. A 10(8): 1173–1181. 13. Kujanpaa, V., Suutala, N., Takalo, T.,and Moisio, T. 1979. Correlation betweensolidification cracking and microstructurein austenitic and austenitic-ferritic stain-less steel weld. Weld. Res. Int. 9(2): 55–76. 14. Brooks, J. A., Thompson, A. W., andWilliams, J. C. 1984. A fundamental studyof the beneficial effects of delta ferrite inreducing weld cracking. Welding Journal63(3): 71-s to 83-s. 15. Brooks, J. A., and Thompson, A. W.1991. Microstructural development andsolidification cracking susceptibility ofaustenitic stainless steel welds. Int. Mater.Rev. 36(1): 16–44. 16. Lippold, J. C., and Kotecki, D. J.2005. Welding metallurgy and weldabilityof stainless steels. Wiley and Sons. 17. Katayama, S., and Matsunawa, A.1984. Soldification microstructure of laserwelded stainless steels. In ICALEO. 18. Katayama, S., and Matsunawa, A.1985. Solidification behaviour and mi-crostructural characteristics of pulsed andcontinuous laser welded stainless steels. InICALEO, pp. 19–25. 19. David, S. A., Vitek, J. M., and Heb-ble, T. L. 1987. Effect of rapid solidificationon stainless steel weld metal microstruc-tures and its implications on the schaefflerdiagram. Welding Journal 66(10): 289-s to300-s. 20. Elmer, J. W., Allen, S. M., and Eagar,T. W. 1989. Microstructural developmentduring solidification of stainless steel al-loys. Metall. Trans. A 20(10): 2117–2131. 21. Vitek, J. M., and David, S. A. 1988.The effect of cooling rate on ferrite in type308 stainless steel weld metal. Welding Jour-nal 67(5): 95-s to 102-s. 22. Nakao, Y., Nishimoto, K., and Zhang,W. P. 1988. Effects of rapid solidification bylaser surface melting on solidification modesand microstructures of stainless steels.Trans. Japan Weld. Soc. 19(2): 20–26. 23. Lippold, J. C. 1994. Solidification be-havior and cracking susceptibility of pulsed-laser welds in austenitic stainless steels.Welding Journal 73(6): 129-s to 139-s. 24. Umeda, T., Okane, T., and Kurz, W.1993. Phase selection during solidificationof peritectic alloys. Acta Mater. 44(10):4209–4216. 25. Lippold, J. C. 1985. Centerline crack-ing in deep penetration electron beam weldsin Type 304L stainless steel. Welding Journal64(5): 127-s to 136-s. 26. Lienert, T. J., and Lippold, J. C. 2003.Improved weldability diagram for pulsedlaser welded austenitic stainless steels. Sci.Technol. Weld. Join 8(2): 1–9. 27. Brooks, J., Robino, C., Headley, T.,and Michael, J. 2003. Weld solidificationand cracking behavior of free-machiningstainless steel. Welding Journal 82(3): 51-sto 64-s.

References

Acknowledgments

....

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Call for Papers JOM­19

19th International Conference on Joining MaterialsHelsingør-Denmark, May 7 to 10, 2017In association with the IIW and supported by The American Welding Society

Your participation at the JOM-19 International event is welcomed. Please fill in a registration form along with a titleand a short abstract of your paper and return it by either e-mail or post before November 2, 2016.

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Main Topics The conference program will cover all aspects of developments in joining and material technology, but papers are espe-cially invited on the following topics:

• Recent developments in joining processes — welding, surfacing, soldering, brazing; • Advances in fabrication techniques; • Developments in joining equipment, including automation and robotics; • Advances in materials, consumables, and weldability; • Applications with relevance to industry needs — automotive, oil and gas, power generation; • New developments in conservation, energy efficiency, and alternative energy resources; • Weld quality, structural properties, and environmental considerations; • Structural integrity and inspection; • Process monitoring, sensors, control; • Mathematical modelling and simulation; • Quality requirements for welds and quality management; • Testing of welds, assessment of defects, and life estimation; • Repair and maintenance of welded structures and plants; • Education, training, qualification, and certification of welding personnel; • Interpretation of International standards for welded fabrications.

28. Arata, Y., Matsuda, F., and Katayama,S. 1977. Solidification crack susceptibility inweld metals of fully austenitic stainlesssteels. Report II. Trans. JWRI. 105–116. 29. Arata, Y., Matsuda, F., Nakagawa,H., and Katayama, S. 1978. Solidificationcrack susceptibility in weld metals of fullyaustenitic stainless steels. Report IV. Trans.JWRI: 21–24. 30. Katayama, S., Fujimoto, T., andMatsunawa, A. 1985. Correlation amongsolidification process, microstructure, mi-crosegregation, and soldification crackingin stainless steel weld metals. Trans. JWRI(7): 123–138. 31. Li, L., and Messler Jr., R. W. 1999.The effects of phosphorus and sulfur onsusceptibility to weld hot cracking inaustenitic stainless steels. Welding Journal78(12): 387-s to 396-s. 32. Kato, T., Fujikura, M., Ishida, K.,Kimura, A., Takeuchi, Y., Kawasaki, N., andYahagi, S. 1981. Effects of chemical com-positions on the properties of austeniticmanganese steels for nonmagnetic applica-tions. Trans. ISIJ 21: 852–861. 33. Reed, R. P. 1989. Nitrogen inaustenitic stainless steels. JOM 41(3):16–21. 34. Brooks, J. A. 1975. Weldability ofhigh N, high MN austenitic stainless steel.

Welding Journal 54(6): 189-s to 195-s. 35. Tate, S. B., and Liu, S. 2014. Solidifi-cation behaviour of laser welded type21Cr–6Ni–9Mn stainless steel. Sci. Technol.Weld. Join. 19(5): 310–317. 36. Goodwin, G. 1987. Development ofa new hot-cracking test — the Sigmajig.Welding Journal 66(2): 33-s to 38-s. 37. Feng, Z., Zacharia, T., and David, S.1997. Thermal stress development in a nick-el based superalloy during weldability test.Welding Journal 76(11): 470-s to 483-s. 38. Ritter, A. M., Henry, M. F., and Sav-age, W. F. 1984. High temperature phasechemistries and solidification mode predic-tion in nitrogen-strengthened austeniticstainless steels. Metall. Trans. A 15(7):1339–1351. 39. Ritter, A. M., and Savage, W. F.1986. Solidification and solidificationcracking in nitrogen-strengthenedaustenitic stainless steels. Metall. Trans. A17(4): 727–737. 40. Suutala, N. 1982. Effect of man-ganese and nitrogen on the solidificationmode in austenitic stainless steel welds.Metall. Trans. A 13(12): 2121–2130. 41. Robino, C., Michael, J., andMaguire, M. 1998. The solidification andwelding metallurgy of galling-resistantstainless steels. Welding Journal 77(11):

446-s to 457-s. 42. Tate, S., and Liu, S. 2014. Fiber laserwelding of high-N, High-MN Austeniticstainless steel. International Congress on Ap-plications of Lasers & Electro-Optics. 603. 43. Hochanadel, P., Lienert, T., Mar-tinez, J., Martinez, R., and Johnson, M.2011. Introduction, weld solidificationcracking in 304 to 304l stainless steel. HotCrack. Phenom. Welds III, pp. 145–160. 44. Honeycombe, J., and Gooch, T.1972. Effect of manganese on cracking andcorrosion behaviour of fully austeniticstainless-steel weld-metals. Met. Constr.4(12): 456–460. 45. David, S. A., and Vitek, J. M. 1989.Correlation between solidification parame-ters and weld microstructures. Int. Mater.Rev. 34(5): 213–245. 46. Blecher, J. J., Palmer, T. A., and De-bRoy, T. 2014. Solidification map of a nick-el-base alloy. Metall. Mater. Trans. A 45(4):2142–2151. 47. Perricone, M. J., and DuPont, J. N.2003. Laser welding of superausteniticstainless steel. Proceedings: Trends in Weld-ing Research, pp. 47–52. 48. Yardy, J., and Laursen, B. 1995.Weld pool microstructures in pulsed-laserwelded austenitic stainless steels. Trends inWelding Research: 165–170.

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