Solidification and Weldability of Nb-Bearing Superalloys

ABSTRACT. The solidification and weldability of experimental superalloys containing systematic variations in Fe, Nb, Si, and C were studied using differential thermal analysis (DTA), microstructural characterization techniques, and Varestraint testing. Microstructural evolution during solidification of DTA samples and gas tungsten arc welds generally occurred by a three-step solidification process. Solidification initiated by a primary L  reaction, which enriches the interdendritic liquid in Nb and C until a eutectic-type L ( + NbC) reaction occurs over a rather broad temperature range. Solidification terminated by a second eutectic-type reaction, L ( + Lave), which occurred over a narrow temperature range. Carbon additions increased the start temperature of the L ( + NbC) reaction and induced a concomitant decrease in the temperature interval of the primary L  stage of solidification. As a result, carbon additions often promoted a significant improvement in solidification cracking resistance. The results from solidification cracks studies are combined with detailed examinations of microstructural morphologies in fusion zone solidification cracks to propose a microstructure morphology classification scheme. The classification scheme provides a phenomenological explanation of the relation between alloy composition, solidification temperature range, microstructural morphology and cracking propensity. This interpretation also leads to an improved correlation between the solidification temperature range and fusion zone cracking susceptibility.

Introduction 

Solidification Cracking 

     The formation of cracks during solidification can occur in ingots, castings, and welds. The phenomenon has been investigated for quite some time and the general features are understood in a qualitative sense (Ref. 1). Work is also in progress to quantitatively determine conditions that lead to solidification cracking (Refs. 2, 3). In general, cracks can form in certain alloy systems during the terminal stages of solidification when a liquid film is distributed along grain boundaries and interdendritic regions. At this stage, shrinkage strains across the partially solidified boundaries can become appreciable. If the terminal liquid is distributed along the boundaries as a continuous film, the strains cannot be accommodated and the boundaries separate to form a crack. In terms of material factors, the solidification temperature range and morphology of the interfacial liquid that exists at the terminal stages of solidification are primary factors which control solidification cracking susceptibility. Solute redistribution plays an important role in solidification cracking as it affects the solidification temperature range and amount of terminal liquid. The effect of the solidification temperature range can be understood in simplified terms by considering its influence on the size of the solid + liquid (mushy) zone.  
     During welding, the mushy region trails behind the liquid weld pool. It is this mushy region that is susceptible to cracking under the influence of restraint. Assuming dendrite tip undercooling is negligible, the interface between the liquid weld pool and mushy zone (i.e., start of the mushy zone) is located where the actual temperature intersects the liquidus temperature of the alloy. Similarly, the interface between the mushy zone and completely solidified weld (i.e., end of the mushy zone) is positioned where the actual temperature intersects the terminal solidus temperature of the alloy. 

Nb-Bearing Superalloys 

     Niobium-strengthened superalloys are used extensively in applications requiring high-temperature strength and resistance to oxidation. Commercial examples include alloys IN625, IN706, IN718, IN903, IN909, and Thermo-Span. These alloys are often used for components that are fabricated by fusion welding. Thus, it is important to establish the relation between alloy composition and cracking tendency in order to avoid fusion zone cracking when these materials are used in critical applications. Previous studies conducted on commercial Nb-bearing superalloys (Refs. 7, 8, 10-12) have sbown that minor variations in Nb, Si, and C have a strong influence on the solidification temperature range, type and amount of secondary phases that form during the terminal stages ofsolidification and result in solidification cracking susceptibility. In general, two types of eutectic-type constituents, /NbC and /Laves, are known to form as the weld metal solidifies. The /Laves constituent is most deleterious in terms of hot cracking as it forms at a lower temperature and thus expands the solidification temperature range. These alloys are also often joined to iron-based alloys such as carbon steels and stainless steels (Refs. 10, 13-15). In these applications,the fusion zone composition can become significantly enriched in Fe and C due to dilution effects, and this composition modification can significantly alter the solidification behavior and associated cracking tendency. For example, Patterson and Milewski evaluated the hot cracking susceptibility of dissimilar metal welds between 304L and IN625 (Ref. 14). The results of that work showed that the diluted fusion zones were more susceptible to cracking than either of the base materials. Subsequent work by Cieslak (Ref. 15) showed that the /Laves constituent forms in the fusion zone of this dissimilar alloy combination. Robitz (Ref. 13) evaluated the solidification cracking of IN625 claddings on C-steel and found that Nb, Si, C, and Fe all increased cracking susceptibility, although no detailed results were reported. DuPont (Ref. 10) showed that solidification cracking in IN625 weld cladding (deposited on C steel) was promoted by formation of the /Laves constituent. Despite the industrial importance of these superalloys and the need to minimize solidification cracking, no general model has been developed for predicting the relationship between alloy composition, solidification behavior (e.g., solidification temperature range and amount of /NbC and /Laves) and resultant weldability. In separate articles, we reported on the development of a solidification model that can be used to predict the solidification temperature range and amounts of /NbC and /Laves constituents as a function of alloy composition (Refs. 16, 17). In this article, the link between solidification behavior and weldability is discussed for a wide range of experimental alloy compositions.

Experimental Alloy Compositions 

     A four factor, two level set of experimental alloys was designed to simulate commercial compositions of interest to this study. The alloy compositions are summarized in Table 1. The alloys exhibit factorial variations in Fe (in exchange for Ni), Nb, Si, and C at two levels. The high- and low-target levels of Nb, Si, and C are set as follows (all values in wt-%): 2< Nb <5, 0.02< C <0.15, and 0.10 < Si < 0.60. These limits were chosen to represent low- and high-composition values of wrought alloys and filler metals as well as composition limits that can arise in fusion welds made between nickel-based alloys and carbon steels. The nominal Cr content was selected at 20 wt-%, which is typically used in many commercial alloys for good corrosion resistance. Several additional alloys with intermediate C contents (Alloys 1.5, 3.5, 7.5, and 11.5) were also investigated. The alloys were prepared at Sandia National Laboratories by investment casting. All samples were melted and poured in vacuum to produce six bars of each composition with approximate dimensions of 170 x 25 x 6.3 mm. A high-Nb alloy (Ni-1) was also added for supplemental differential thermal analysis measurements. Small cast buttons of this alloy were produced from high-purity powders (> 99.99 wt-%) by gas tungsten arc melting. 

Differential Thermal Analysis 

     Differential thermal analysis (DTA) was conducted on a Netzsch STA 409 differential thermal analyzer using 500 to 550 mg samples. The DTA system was calibrated within 2°C of the melting point of pure Ni. Samples were melted and solidified under flowing argon in alumina crucibles using pure Ni as the reference material. The liquidus temperatures were determined by heating samples slowly at a rate of 5°C/min to approximately 10°C above their liquidus temperature.

Varestraint Testing 

     The solidification cracking susceptibility of each alloy was evaluated using the Varestraint test (Refs. 18, 19). The as-cast samples were machined to typical subsize Varestraint specimens (165 x 25 x 3.2 mm). The welds were produced under the following parameters: 95 A, 10 V (2.5 mm arc length), and 3.3 mm/s travel speed with high-purity argon shielding. A 3.2-mm diameter W-2ThO electrode was used with a 60-deg tip angle. In order to acquire adequate statistics, three tests were conducted on each alloy at an augmented strain of 2.5% to simulate highly restrained welds. The maximum crack length (MCL) was used as the indicator of cracking susceptibility and was measured using a light optical microscope at 100X magnification. The level of 2.5% augmented strain was selected on the basis of a large number or tests at varying strain levels (conducted using essentially identical testing equipment, welding conditions and weld sizes) on a range of alloys with similar solidification characteristics and solidification microstructure (Refs. 8, 20, 26). These results have shown that these test conditions, for alloys that solidify over similar temperature ranges and which exhibit MCL values spanning the range observed in the current study, the MCL is generally saturated at the 2.5% strain level. At these strain levels, it is normally expected that cracking will span a major fraction of, if not the entire, length of the mushy zone. 

Microstructural Characterization 

     Light optical microscopy (LOM) was conducted on Varestraint and DTA samples polished through 0.04 µm colloidal silica and electrolytically etched in a 10% chromic acid +90% water solution at 3 V. Scanning electron microscopy was conducted on weld metals (prepared using the same procedure) using a JEOL 6300 field emission gun scanning electron microscope (FEG-SEM) at an accelerating voltage of 15kV. The Varestraint samples were mounted in thermal setting epoxy to provide a planar view of the solidification cracks which intersected the specimen surface. After mounting, the surfaces were lightly ground with 600-grit SiC paper to remove a very thin layer of surface material and provide a flat area in the crack region for subsequent polishing and etching. Quantitative image analysis (QIA) was conducted on autogenous GTA welds at locations far removed from solidification cracks. Area fractions of total eutectic-type contents and, where possible, individual /NbC and /Laves eutectic-type constituents were measured along the centerline of each weld with at least ten SEM photomicrographs. Area fractions were assumed to be equivalent to volume fractions. Electron probe microanalysis (EPMA-WDS) was conducted on a JEOL 733 probe at an accelerating voltage of 15kV and beam current of 20 nA. All EPMA samples were mounted in thermal setting epoxy, polished flat to a 0.3-µm finish using an alumina slurry, ultrasonically cleaned in acetone and carbon-coated prior to analysis. The K lines were used for Fe, Ni, Cr, and Si, while the L line was used for Nb. Raw data were reduced to weight percentages using a ZAF algorithm. 
 
 

Solidification 

     The solidification behavior of these alloys has been described through microstructual analysis (Ref. 16) and solidification modeling (Ref. 17) in separate articles. The important features are summarized below as they provide a necessary framework for detailed interpretations of the weldability results. Typical DTA cooling scans for Alloys 2 and 16, which represent the two types of reaction sequences observed among the experimental alloys, are presented in Fig. 1. The corresponding microstructures of GTA welds and DTA samples are presented in Fig. 2. The first type of solidification sequence, which was observed for all alloys except 2, 3.5, and 4, can be described by a three-step process: 1) Primary L  solidification in which the interdendritic liquid becomes enriched in Nb and C; followed by 2) a eutectic-type L  (+NbC) reaction, which depletes the interdendritic liquid of C; and 3) termination of solidification by a second eutectic-type reaction L  ( + Laves). The broad peak associated with the L  ( + NbC) reaction indicates this transformation occurs over a wide temperature range, while the narrow peak associated with the L  ( + Laves) reaction indicates this step occurs over a small temperature interval (Ref. 16). Identification of the NbC and Laves phases shown in Fig. 2 was confirmed through EPMA measurements on coarse phases in DTA samples and backscattered electron kikuchi patterns collected from finer phases in weld metal samples (Refs. 16, 21). The Ni-based alloys with high C/low Nb (Alloys 2, 3.5, and 4) represent the only exception to this general solidification sequence. These alloys did not show any exothermic peaks after the L  ( + NbC) reaction and did not exhibit any of the Laves phase in the as-solidified micro structure. This indicates that solidification terminated with the L  ( + NbC) reaction. The amounts of total eutectic-type constituents (/NbC + /Laves) and, where possible, individual /NbC and /Laves constituents, are summirized in Figs. 3 and 4. The data are arranged to permit comparisons among Ni-based and Fe-based alloys with similar levels of solute elements (Nb, Si, and C). 
     As discussed in detail later, the solidification cracking resistance of these alloys is controlled, to a large degree, by the relative amounts of /NbC and /Laves constituents. Thus, it is instructive to consider the relation between alloy composition and relative phase fractions of /NbC and /Laves displayed in Figs. 3 and 4. At similar levels of minor alloy additions (Nb, Si, and C), the iron-based alloys generally form higher levels of total eutectic-type constituents. Measurements of the individual eutectic-type constituents (Fig. 4) indicates that this response can be attributed primarily to the formation of higher amounts of /Laves in the Fe-based alloys. The addition of Fe to Nb-bearing Ni-based alloys is well known to stabilize the Laves phase (Ref. 22). Thus, Fe additions to the austenite matrix, at the expense of Ni, can be expected to promote the formation of more /Laves, as observed here. The addition of Si has a similar effect, as the amount of /Laves always increases with increasing Si within a given matrix composition (in Fig. 3, compare Alloys 5 to 7, 6 to 8, 13 to 15, and 14 to 16). The Laves promoting tendency of Si has also been documented in superalloys (Ref. 23). Niobium additions promote higher amounts of total eutectic. This is to be expected, since each of the secondary phases that form in the eutectic-type structures (NbC and Laves) are both highly enriched in Nb. At high C levels, Nb additions generally promote more /NbC (compare Alloys 2 to 6, 4 to 8, 10 to 14, 11.5 to 16). When the C and Fe levels are both high, Nb additions will increase the amounts of both the /NbC and /Laves constituents (compare Alloys 10 to 14 and 11.5 to 16). Alloys with high C levels (even-numbered alloys) form large amounts of the /NbC eutectic-type constituent. The presence of NbC and Laves has been observed in a large number of commercial superalloys, including Alloys IN625 (Ref. 12), IN718 (Ref. 8), IN909 (Ref. 24), IN903 (Ref. 23), IN519 (Ref. 11), Thermo-Span (Ref. 26), and IN625 weld cladding (Refs. 10, 27). The experimental alloys used in this work properly simulate the wide range of microstructures formed in commercial superalloys and weld cladding. 

Fig. 2 -- Microstructures of DTA samples and GTA welds for Alloy 2 and Alloy 16. A -- Alloy 2 DTA sample; B -- Alloy 2 weld; C -- Alloy 16 DTA sample; D -- Alloy 16 weld.  
 

Fig. 4 -- Quantitative image analysis results showing amount and individual /NbC and /Laves constituents.  

     The solidification reactions in these experimental multicomponent alloys and commercial alloys are similar to those expected in the pure ternary Ni-Nb-C system (Refs. 28, 29). The liquidus projection for this system is shown in Fig. 5. The liquidus projection exhibits three primary phase fields that are of interest here, , NbC, and Ni₃Nb. A primary C (graphite) phase field exists at high C contents, which is not of importance. As noted earlier, additions of Fe, Cr, and Si to the Ni-Nb system are well known to promote Laves at the expense of Ni₃Nb in commercial superalloys, as well as the experimental alloys utilized in this work. Thus, by replacing Ni₃Nb with Laves, the Ni-Nb-C liquidus projection can be utilized as a guide in developing a description of the solidification reactions in these alloys (Ref. 29). 
     Solidification begins with formation of primary  dendrites which, upon forming, reject Nb and C to the liquid (Ref. 16). As solidification proceeds, the liquid composition moves away from the Ni corner, becoming progressively richer in Nb and C until the line of two-fold saturation between  and NbC is reached. At this point, the maximum solubility of Nb and C in the austenite matrix is exceeded and the  and NbC phases form simultaneously from the liquid by a eutectic-type reaction as the liquid composition travels down the line of two-fold saturation. Due to its high C content, the formation of NbC depletes the liquid of C while the Nb content of the liquid continues to increase. For Alloys 2, 3.5, and 4, solidification is completed along the line of two-fold saturation between  and NbC, as no Laves phase was observed in these alloys. The initial reaction sequence is similar for the remaining alloys, except that solidification does not terminate with the L  ( + NbC) reaction. Instead, the liquid composition continues to travel down the line of two-fold saturation between  and NbC, becoming progressively enriched in Nb (and depleted in C).

Fig. 5 -- Liquidus projection for the Ni-Nb-C system. 
 

Fig. 6 -- A -- Calculated primary solidification paths of Alloys 1 through 4 superimposed on the experimentally determined liquidus projection; B -- calculated primary solidification paths of Alloys 2 and 10 superimposed on the experimentally determined liquidus projection (Refs. 16, 17). 

     The effect of matrix composition is presented in Fig. 6B. Alloys 2 and 10 have very similar levels of solute content, but different matrix compositions (Alloy 2 is Ni-based while Alloy 10 is Fe-based). The matrix composition has a strong influence on the segregation potential of Nb (Ref. 16) as measured through the equilibrium distribution coefficient, k_Nb. The value of k_Nb in the Ni-based alloys is equal to 0.46, while k_Nb in the Fe-based alloys is only 0.25 (Ref. 16). As k_Nb is reduced, Nb segregates more aggessively to the liquid. Thus, at any given level of C in the liquid, the Fe-based alloys will always possess more Nb. As a result, the Fe-based alloys will have more liquid remaining after primary solidification, and the Nb content in that remaining liquid will be higher. These effects from the reduced k_Nb value are the primary reasons for the higher amount of total eutectic-type constituents observed in the Fe-based alloys -- Fig. 3. The validity of the solidification model was confirmed by comparing the calculated and measured amounts of /NbC and /Laves constituents in the experimental alloys, and reasonable agreement was found (Ref. 17).

Weldability 

Effect of Solidification Temperature Range 

     Figure 7 shows the maximum crack length (MCL) of each alloy. (The alloy compositions are summarized in Table 1). Within the Ni-based alloys, there is a clear separation in weldability among the low-C alloys (0.017 wt-% C) with relatively poor weldability and the high-C alloys (=0.052 wt-% C), which show very good resistance to solidification cracking. Within the Fe-based alloys, the addition of C is only beneficial when the Nb content is low (1.93 wt-%), and the C level must be above 0.10 wt-% to provide the advantageous effect. For example, Alloys 11.5 and 12 have essentially identical levels of all other elements except C (Alloy 12 has 0.079 wt-% C and Alloy 11.5 has 0.116 wt-% C). This small variation in C content leads to a substantial difference in the MCL values. Within the Fe-based alloys with high Nb (Alloys 13-16), C has no beneficial effect even at the 0.21 wt-% level. 
 

Fig. 7 -- Maximum crack length of each alloy. 

     Figure 8 plots the MCL as a function of the solidification temperature range, and no correlation is readily apparent. Many alloys exhibit a rather wide solidification temperature range, but show good weldability (low MCL). It should be emphasized that Cieslak (Ref. 7) was also unable to obtain a correlation between MCL and T_s in his previous weldability reports on alloy IN625. This lack of correlation between MCL and T_s suggests that crack propagation is being affected by the amount and distribution of solid and liquid phases within the mushy zone. Such potential effects can be assessed by conducting detailed microstructural characterization of the eutectic-type constituents in the vicinity of solidification cracks. These constituents transform from the terminal liquid, which existed in the grain boundary and interdendritic regions just before solidification is complete. Thus, the amount and morphology of these constituents provides a basis for developing an understanding of the distribution of solid and liquid phases within the mushy zone.

Table 3 - /Laves Average Constituent Compsitions.

Alloy/Sample	Fe	Ni	Cr	Nb	Si
Alloy 7 DTA	4.9  0.2	59.4  0.8	10.2  1.0	24.3  0.7	1.7  0.1
Alloy 7.5 DTA	4.9  0.2	58.1  1.7	11.4  0.8	23.5  1.3	2.0  0.1
Alloy 8 DTA	4.4  0.2	60.3  0.5	11.3  0.3	21.5  0.5	2.5  0.2
Ni-1 GTA melt	7.0  0.1	52.2  0.4	16.7  0.2	23.1  0.5	1.0  0.1
Alloy 13 Weld	31.4  1.8	29.8  1.1	16.6  0.6	21.1  1.3	0.3  0.1
Alloy 13 DTA	31.2  0.9	30.4  0.5	17.3  0.2	20.9  1.0	0.3  0.1
Alloy 14 Weld	31.0  1.1	29.5  0.8	16.8  0.7	22.4  2.4	0.3  0.1
Alloy 15 Weld	32.7  0.5	29.7  1.0	16.7  0.4	20.1  0.9	0.9  0.1
Alloy 16 Weld	32.7  0.4	30.8  0.7	16.9  0.1	18.3  0.7	1.3  0.1
Alloy 16 DTA	28.6  1.9	33.4  1.4	16.5  0.9	19.8  1.2	1.7  0.1

Microstructural Evolution and Solidification Cracking 

     The general solidification process of these alloys was described in the solidification section above. By combining the microstructural characterization results of Varestraint samples with the solidification sequences already discussed, it is possible to extend this general solidification description to identify four individual types of microstructural morphologies and their effect on cracking susceptibility. This proposed classification scheme sheds light on a more appropriate relation between MCL and T_s which could not be resolved by the initial plot in Fig. 8. All of the alloys can be grouped into one of four morphological classification schemes, and Table 4 lists the classification scheme that applies to each alloy. 

Fig. 13 -- SEM photomicrograph showing the /NbC and /Laves morphology in the solidification crack of Alloy 16, Type III microstructure.

Fig. 14 -- SEM photomicrograph showing the microstructural morphology in the solidification crack of Alloy 3, Type IV microstructure.

     The addition of C is beneficial as it increases the start temperature of the L  ( + NbC) reaction and reduces the temperature interval of the primary L  stage of solidification. This favors formation of Types I and II microstructures, which show very good weldability. This beneficial effect can be lost if more than approximately 2.3 vol-% of the /Laves constituent forms, at which point the Type III microstructure develops with a concomitant reduction in cracking resistance. The /Laves constituent is promoted by additions of Fe, Nb, and Si. In the Ni-based alloys that have less than 11 wt-% Fe, the segregation potential of Nb is relatively low and the /Laves eutectic composition is relatively high. These two effects work together to reduce the amount of /Laves that forms, and C additions to the Ni-based alloys are consistently beneficial. Within the Ni-based alloys, the transition from Type I to Type II microstructures occurs when the Nb content is increased. The Type III microstructure does not form in the Ni-based alloys since the /Laves constituent is always below 2.3 vol-%. 
     When the Fe content is high, the segregation potential of Nb is increased and the /Laves eutectic composition is reduced. Thus, at equivalent levels of solute elements (Nb, Si, and C), the Fe-based alloys will always form more /Laves than the Ni-based alloys -- Fig. 4. As a result, the Fe-based alloys never form the Type I microstructure, and the Type II microstructure can only develop in the low-Nb alloys (Alloys 10 and 11.5). The sensitivity of microstructure type and resultant weldability to C content is reflected in the differences displayed by Alloys 11.5 and 12. Both have similar levels of Nb and Si, but Alloy 11.5 has slightly higher C (0.116 wt-% vs. 0.079 wt-%). At this level of Fe, Nb, and Si, this small difference in C content is enough to induce a change in microstructure type.  
 
 

These two microstructural morphologies are compared in Fig. 16. When the Fe-based levels have high Nb, a high amount of /Laves forms and leads to the Type III microstructure. Finally, for the low-carbon alloys, the solidification path always travels along the /Nb binary side of the liquidus surface and leads to the Type IV microstructure, which have relatively poor weldability. These relations were discussed qualitatively in this article. Results of solidification modeling presented in previous articles can be used to predict these interrelations in a quantitative sense (Refs. 16, 17).

Conclusions 

     The solidification and weldability ofexperimental Nb-bearing superalloys with systematic variations in Fe, Nb, Si, and C was investigated by differential thermal analysis, Varestraint testing, and  
microstructural characterization techniques. The following conclusions can be drawn from this work: 
     1) The general solidification sequence of these alloys can be described by a three-step process: 1) Primary L  solidification in which the interdendritic liquid becomes enriched in Nb and C; followed by 2) a eutectic-type L  ( + NbC) reaction, which depletes the interdendritic liquid of C; and 3) termination of solidification by a second eutectic-type reaction L  ( + Laves). This solidification reaction sequence is similar to that expected in the ternary Ni-Nb-C system and observed experimentally in many commercial alloy systems. 
     2) Carbon additions have a strong influence on the temperature interval of the primary L   stage of solidification. As the carbon content is increased, the L  ( + NbC) reaction temperature increases and, as a result, reduces the temperature interval of the primary L  stage of solidification. In the Ni-based alloys, C additions above 0.052 wt-% significantly improve the solidification cracking resistance. Carbon additions to the Fe-based alloys are only advantageous when the Nb level is below 2 wt-%, and the C level must be above 0.12 wt-% for the beneficial effect to be observed. 
     3) There is generally no direct relation between the actual solidification temperature range and fusion zone cracking susceptibility of these experimental alloys, suggesting that crack propagation through the mushy zone is affected by the amount and morphology of solid and liquid phases in the mushy zone. 
     4) A classification scheme is proposed to provide a phenomenological explanation of the relation between the solidification temperature range, microstructural morphology, and cracking propensity. This morphological description leads to an improved correlation between the (effective) solidification temperature range and fusion zone cracking tendency. 
 

     The authors extend their thanks to Dr. M. J. Cieslak at Sandia National Laboratories for usefull discussions and review of the manuscript. One author (JND) gratefully acknowledges financial support for this research from the American Welding Society Fellowship Award. Preparation of the experimental alloys by B. Damkroger and M. Maguire at Sandia National Laboratory is also greatly appreciated. Sandia is a multiprogram laboratory operated by Sandia Corporation, a Lockheed Martin Company, for the United States Department of Energy under Contract DE-ACO4-94AL85000. 

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