The Welding of Nickel Superalloys: Cracking Mechanisms and Process Mapping Due to the localised heating/melting of material, laser processing (both SLM and DLF) can be

considered analogous to welding. It is well documented that precipitation hardenable nickel superalloys with a high  fraction are susceptible to weld cracking. Figure 2.45 shows a plot presented by Donachie [2] illustrating the Al and Ti ( forming elements) and an empirical line representing the boundary between materials considered to be weldable and un-weldable due to their cracking susceptibility. This method of weldability assessment is also supported by Henderson et al. [107] although it must be emphasised that this does not take into account the mechanisms or the complex microstructural effects governing weldability; the values for CM247, CMSX486 and IN625 have been added to the plot. Note that CM247 and CMSX-486 lie firmly in the ‘un-weldable’ region whereas IN625 appears in the weldable portion of the plot.

Figure 2.45: Plot (Reconstructed from source) showing the Al and Ti contents of various nickel superalloys (including CM247LC, CMSX-486 and IN625). The dashed line indicates an empirical boundary between

weldable and un-weldable material due to cracking susceptibility [2].

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Work by Dye et al. [108] has shown it is possible to construct a map for welding of Ni-superalloys by identifying the potential defects which may occur as a function of the effective weld power and weld speed. An example of this type of weldability map (developed for IN718) is shown in Figure 2.46.

Figure 2.46: Typical (TIG) weldability map for IN718 [108]

The research by Rush et al. [109] also presents a parametric study aimed to map the process parameters of laser welding to cracking within Rene 80. The research presented by Egbewande et al. [110] and Zhong et al. [99] (both regarding IN738) are less complete, but still show some attempt to relate to the process parameters to the crack occurrence within the material. In general terms all three of these studies show that by reducing the heat input the occurrence of cracking is reduced.

Based on a review of the relevant literature, four cracking mechanisms associated with the welding of nickel superalloys have been identified; these mechanisms are identified, discussed and documented relationships to processing parameters identified:

69 i. Solidification Cracking (Hot Tearing)

This mechanism is described in general by traditional texts on the subject such as that of Kou [111] or DuPont et al. [112] as occurring in the partially liquid phase during the solidification of the material. Dye et al. [108] present an extensive numerical analysis of weld cracking in nickel superalloys and describe solidification cracking succinctly as occurring when the solid fraction (fs) is in the range of 0.7-0.9 (depending on material and sources). During this range, the movement of the remaining liquid is restricted by the growing dendrites and, under the action of the solidification stress, it is unable to backfill into the contracting regions causing cracks to open. Kou [111] expands on this by stating that the solidification cracking susceptibility is increased with increasing solidification range (i.e. Difference between solidus and liquidus of material). It is also stated that the crack wall shows a distinctive dendritic morphology where the crack has formed between the solidifying dendrites; Figure 2.47 shows a micrograph presented by Kou [111] of this phenomenon in a steel solidification crack.

Similar observations of an “unusual cellular structure” on the crack wall of thin walled castings of the nickel superalloy MAR-M246 were reported by Winstone & Northwood [113].

Figure 2.47: Micrograph showing dendritic structure on the wall of a solidification crack present in 308 stainless steel [111].

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The weld mapping performed by Dye et al. [108], regarding IN718, shows that solidification cracking tends to occur under low welding speeds and high effective power (i.e. high heat input; Figure 2.46). The research presented by Rush et al. [109] identifies the occurrence of solidification cracking to be restricted to cracks showing a very high Average Crack Length (ACL) occurring under high power/small beam diameter (i.e. high heat input) laser welds of Rene 80. Egbewande et al. [110] also supports the idea of solidification cracking under high heat input (in laser welded IN 738) as a significant reduction in cracking is shown when moving from 0.5 – 1 m/min weld speed. This reduction in cracking is linked to the elimination of cracks which appear to be caused during solidification (based on the micrographs presented).

ii. Liquation Cracking

Grain boundary liquation in nickel superalloys is widely reported; The papers presented by Henderson et al. [107] and Attallah et al.[114] summarise this phenomenon by describing it as occurring when material is heated rapidly to a high temperature, but not above the overall melting point of the bulk material. Under these conditions some grain boundary phases are unable to dissolve rapidly enough into solid solution and can form low-melting point phases (such as eutectics) at the grain boundaries. These phases can melt and form liquid grain boundary films which can act as crack initiation points under residual or thermal stresses.

The precise composition of the liquid phases varies from alloy to alloy. In solid-solution and

-strengthened alloys, liquation is generally attributed to carbides [109, 115, 116], however in -strengthened, grain boundary  (or the formation of a / eutectic under rapid heating) has been attributed to liquation [99, 110, 117].

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Figure 2.48 shows a micrograph presented by Thompson & Radhakrishnan [116] revealing the interdendritic liquation in a sample of IN718 following heating and quenching. In this situation stress was not applied so the liquated phase resolidified in the interdendritic region without cracking. Figure 2.49 shows a micrograph presented by Sidhu & Chaturvedi [117] of the HAZ (Heat Affected Zone) in welded Rene80; here the thermal stress imposed by the weld has shown that the liquation film acts as a crack initiation point along the grain boundary. The propagation of cracks along grain boundaries is further explored by Zhong et al. [99] in laser deposited IN738 onto a DS superalloy substrate. The micrograph in Figure 2.50 shows the liquation occurring in the HAZ of the substrate and the crack propagating along a grain boundary into the deposited IN738.

Figure 2.48: Micrograph showing interdendritic liquation of heated and quenched IN718 [116].

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Figure 2.49: SEM showing grain boundary liquation and crack propagation in the HAZ of Rene 80 [117].

Figure 2.50: Micrograph showing laser deposited IN738 on a DS superalloy substrate. The liquation occurs in the HAZ of the substrate and the crack formed propagates along a grain boundary into the

deposited material [99].

The micrograph presented by Odabaşi et al. [118] in Figure 2.51 confirms the region subject to liquation to be in the HAZ rather than in the fusion zone (laser-welding of IN718). The research also shows that liquation is more distinct in samples welded under lower energy conditions (i.e. faster weld speed) rather than the higher energy input conditions. This is supported by the weld map for IN718 presented by Dye et al.[108] (Figure 2.46) showing liquation cracks to occur under low power and high speed welding conditions.

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Figure 2.51: Micrograph showing the occurrence of liquation in the HAZ away from the fusion zone in IN718 [118].

iii. Ductility-Dip Cracking (DDC)

DDC is not clearly separated in the literature and is often included in the general categories of

‘reheat cracking’ and ‘hot cracking’; but recently research by Dupont et al. [112] has clarified DDC as a separate mechanism. Lippold and co-workers define it as a cracking caused by the reduction in ductility occurring in nickel alloys at intermediate temperatures (0.4-0.7 Tm) [112, 119]. Much of the literature referring to DDC is in relation to weld filler materials;

however research by Kim et al. [50] shows a ‘ductility-dip’ in tensile data presented for CM247LC in the range 700C - 900C.

Two differing microstructural mechanisms have been suggested for the occurrence of DDC within the literature: That suggested by Lippold et al. and Young et al.

Lippold et al. [119-123] suggest DDC to be a ‘creep-like’ mechanism which operates at a high enough temperature to promote grain boundary sliding, but below the threshold of dynamic recrystallisation [119, 122]. It is concluded that the grain boundary sliding leads to stress concentrations and ultimately the formation of voids on grain boundary features. Within these, papers by Lippold et al [119, 120, 122] discuss the factors influencing DDC occurrence. Triple point boundary intersections were shown to be points of severe stress

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concentrations in material showing relatively smooth/straight grain boundaries. Fine grain boundary particles (carbides) were able to provide some grain boundary locking in order to reduce the occurrence of voids at the triple point intersections; however the particles themselves became sources of stress concentration. Tortuous grain-boundaries can also provide greater resistance to DDC when combined with the effect of fine grain boundary particles and ultimately larger grain boundary particles added a microstructural locking effect further adding to the DDC resistance. Figure 2.52 presented by Lippold & Ramirez [122]

shows the general scheme from ‘triple point stress’ to ‘particle stress’ to the division of stress between the particles and the ‘tortuous grain boundaries’.

Figure 2.52: Diagram showing the high strain concentrations at triple point boundaries (a) the movement of the high strain regions with the addition of grain boundary particles (b) and finally the division of the

high the inclusion of a tortuous grain boundary [122].

The final point to note is that Lippold et al. [120] also concluded that the susceptibility to DDC is increased at high angle grain boundaries.

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Research presented by Young et al. [124] suggests a different mechanism that is due to the precipitation of partially coherent carbides at grain boundaries under intermediate temperatures; the stress imposed by these carbides acts upon the grain boundaries forming microscopic voids which act as crack initiation points under larger stresses. This is shown schematically in Figure 2.53.

Regardless of the formation mechanism, the onset of DDC can be characteristically seen as a series of voids lying on a grain-boundary between carbide particles, Figure 2.54.

Figure 2.53: Diagram (Reconstructed from source) showing schematically the proposed mechanism for DDC by the formation of voids from localised stresses around carbides precipitated at grain boundaries

[124].

Figure 2.54: SEM micrograph (poor quality original) showing DDC void formation along a line of carbides [124].

iv. Strain-Age Cracking (SAC) or Post-Weld Heat Treatment (PWHT) Cracking SAC is not typically associated with the welding process, but is associated with a post-weld heat-treatment (PWHT) of a  precipitation hardenable Ni-base superalloy and is well

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documented. SAC may occur during stress-relief treatment following a weld, during an ageing treatment or in the ramp phase of a solution treatment. When heated to within the ageing region of the alloy, two competing actions occur; the precipitation of the  phase and the much slower relaxation of residual stress due to welding. The precipitation of  increases the material strength whilst reducing ductility [2, 107]). Rowe [125] reports that an additional stress is induced due to the reduction in volume inherent in the precipitation of  from solid solution. The combination of these two stresses combined with the reduction in ductility can exceed the material strength causing cracks to open. Henderson et al. [107] reports that these cracks typically open at grain boundaries with particulate carbides acting as initiation points.

Figure 2.55 presented by Donachie [2] graphically shows this sequence of events:

Figure 2.55: Diagram showing the potential sequence of events leading to SAC [2].