Comparative study on dissimilar welding of Inconel 718 and Udimet 500 precipitation strengthened nickel based superalloys

Abstract

Microstructural, mechanical and weldability aspects in the similar and dissimilar welds of alloy 718 and alloy 500 nickel based superalloys have been investigated. Alloy 500 weld metal showed high tendency of titanium to the segregation. Coalescence of the microvoids led to propagation of hot solidification microfissures. The alloy 718 weld metal displayed the formation of Nb rich low melting eutectic type morphologies, which can reduce the weldability. The microstructure of dissimilar weld metal with dilution of 65 wt- displayed semideveloped dendritic boundaries. The less segregation and decrease in the low melting eutectics caused less susceptibility of dissimilar weld to solidification cracking. The segregation elimination phenomenon has occurred in the heat affected zone of alloy 500. In the partially melted zone, remelted and resolidified regions have been observed. These locations provided sites for nucleation of liquation cracks. For the alloy 718 heat affected zone, dissolution of γ″-Ni₃Nb needle-like precipitations has taken place. It was the chief reason for sharp decline of the microhardness. The heat affected zone of alloy 500 revealed intense liquation cracking, in which the crack is initiated at the partially melted zone. The hot liquation cracking in the heat affected zone of Alloy 718 was observed as a result of γ″-Ni₃Nb dissolution.

Keywords

Introduction

Precipitation strengthened nickel based superalloys are extensively utilised in most of the critical industrial applications such as power plants, land based and jet gas turbines, aerospace, oil refineries, chemical processing and military equipment as an irreplaceable selection.^1–8

Although their excellent mechanical properties and resistance to high temperature corrosion make them to be used widespread, welding and weldability of precipitation strengthened nickel based superalloys have caused challenging issue for the metallurgist and metallic material researchers.^9–12 The precipitation strengthened nickel based superalloys are divided into two general groups. In the first, in which the strengthening agent is dispersed in intermetallic precipitations of gamma prime (γ′) with chemical composition of Ni₃(Ti,Al), the weldments are prone to strain age cracking.^13–16 In addition, in the more recent studies, it is discovered that reaction of TiC precipitations with the austenitic matrix, remelting of γ–γ′ eutectic structures and reaction of gamma prime particles with the austenite have resulted in the formation of liquid films and liquation cracking in the weld heat affected zone (HAZ).^17,18 Moreover, in some investigations, it is reported that relatively strong segregation of Ti into the interdendritic regions of the weld metal during the solidification process has led to the broadening of the brittleness temperature range and increasing of the susceptibility to solidification cracking.^19,20 The sensitivity of welds in the γ′ strengthened nickel based superalloys to the strain age cracking, liquation and solidification cracking is so harmful that the title of ‘unweldable’ is proposed for them.¹³

In the newer generation of precipitation strengthened nickel based superalloys, i.e. γ″ strengthened nickel based superalloys, niobium has been used as the strengthening agent instead of Ti and Al. Nb, by forming the dispersed intermetallic precipitations, called gamma double prime (γ″) with chemical composition of Ni₃Nb in the austenite binder, can greatly overcome the sensitivity of these alloy to strain age cracking, but unfortunately, the susceptibility of the weld to the solidification and liquation cracking has increased.^21–23 The researches demonstrate that Nb, which segregates into the interdendritic boundaries, significantly decreases the solidus temperature of the weld metal, resulting in the higher solidification cracking susceptibility.^24–26 On the other hand, niobium can participate in the formation of low melting eutectic constituents such as γ-NbC and γ-Laves, which are formed at the terminal solidification of weld metal. These phases can lead to more reduction in the solidus temperature of the weld, and thus, the weld indicates more severe susceptibility to solidification cracking. In addition, the HAZ of the weld contains some types of precipitations such as NbC, Laves, Ni₃Nb and δ, which react with the surrounding austenite matrix, leading to formation of liquid regions and liquation cracking.^27,28

The studies and investigations performed in the spectrum of welding science and engineering of precipitation strengthened nickel based superalloys have generally accentuated on the microstructural evolutions,^29–31 mechanical characteristics,^32–37 weldability behaviour and compositional alterations^13,18,23,38 of γ′ strengthened nickel based superalloys or γ″ strengthened nickel based superalloys welds, individually. Negligible scientific research studies have been fulfilled on the dissimilar welding of γ′ strengthened nickel based superalloys to γ″ strengthened nickel based superalloys. In the light of discussion above, the current research is organised to emphasise the comparative evaluation of microstructural, mechanical and weldability aspects of dissimilar welding of γ′ strengthened nickel based superalloys to γ″ strengthened nickel based superalloys.

The research includes similar and dissimilar welding of γ′ and γ″ strengthened nickel based superalloys, in which the microstructural characterisations of weld metals, interfaces, HAZs and base metals, and the dependent microstructural features such as segregation, type of precipitations, morphologies, distribution, phase chemical composition analysis and their influences on the mechanical and weldability properties have been surveyed by details.

Experimental

To perform the similar and dissimilar welding of γ′ strengthened nickel based superalloys to γ″ strengthened nickel based superalloys, the following materials were selected: the first contains high content of Nb and the second contains high content of Ti/Al. These were 12 mm thick plates of alloy 718 (i.e. Inconel 718) precipitation strengthened nickel based superalloy strengthened by γ″-Ni₃Nb strengtheners received in hot forged condition, and alloy 500 (i.e. Udimet 500) precipitation strengthened nickel based superalloy strengthened by γ′-Ni₃(Ti,Al) strengtheners received in the as cast condition.

The nominal chemical compositions of the base materials are given in Table 1. As known, autogenous tungsten inert gas welding is the best fusion welding process for basic and preliminary scientific and technological study. Thus, the welding operations were performed without using filler materials and electrodes (i.e. autogenous welding) utilising gas tungsten arc (GTA) welding process. The alloy 718 and alloy 500 base metals were individually remelted by single run autogenous GTA welding to produce bead on plate welds. Dissimilar base metals were also welded to each other using single run autogenous GTA welding by employing a square butt weld joint to obtain different dilution level. Dilution of 65 (i.e. in this case, the weld metal consists of 35 wt- alloy 718 and 65 wt- alloy 500) was one of the best dilutions obtained, which is presented in the current research. This dilution level is reproducible when the fixed and appropriate weld joint design and welding parameters are used.

Table 1

Nominal chemical compositions of base materials

	Chemical composition/wt-
Type of base materials	Elements
Type of base materials	C	Ni	Cr	Co	Fe	Mo	Nb	Ti	Al
Alloy 718 base metal	0·01	Rem.	18·5	…	22	3	5	1·1	0·3
Alloy 500 base metal	0·01	Rem.	20·5	19	…	3	…	2·7	1·9

In addition, the welding voltage, current and speed were simultaneously recorded during the welding. The welding parameters are as follows: current = 300 A, voltage = 12 V and travel speed = 4·3 mm s⁻¹. The high value of welding current is for ensuring sufficient penetration and weld metal appropriate mixing. Three welds were made and used for each weld type for reproducibility. The results presented in the present research are observed in all the welds.

After the welding operations, abrasive cutting and cleaning, the specimens were prepared according to the standard metallographic procedures. Marble etching solution with chemical composition of 5 g CuSO₄+100 mL HCl+100 mL H₂O was used for revealing the weldment microstructures. The identification of the microstructural characteristics was carried out by a Leitz Wetzlar Aristomet optical microscope equipped with Nikon ACT Version 2.70 software. Further microstructural investigations were performed by a Zeiss EVO 50XVP scanning electron microscope (SEM) with accelerating voltage of 30 kV equipped with Oxford Instrument 7060 energy dispersive X-ray spectrometer (EDS) for spot, line and map weight analysis.

For measurements of mechanical properties, i.e. hardness, the Vickers microhardness tester of Future-Tech Corporation FM-700 model, with indentation weight of 1000 gf and indentation time of 15 s, was utilised. The measurements were continuously fulfilled at 500 μm steps starting from the alloy 500 base metal, then across the weld metal and ending in the alloy 718 side. The results obtained were drawn as Vickers microhardness values versus the distance.

Results and discussion

Microstructural characterisation of base metals

Figure 1a shows the microstructure of alloy 500 base metal in as cast condition. Fully developed and coarse dendritic structure accompanied with severe alloying segregation in the interdendritic regions can be seen. In this figure, dendritic cores have appeared in light colour and the interdendritic regions containing segregation show more darkness in comparison to the cores. In addition, in the most interdendritic locations, black chains are distinguished exhibiting probable formation of secondary phases, which can be either the precipitations or eutectic structures, resulting in the solidification process. Figure 1b has illustrated the electron microphotograph of alloy 500 microstructure at higher magnification. The austenite matrix has thoroughly been occupied by the cubic precipitations of γ′-Ni₃(Ti,Al). These dispersed particles are responsible for the strengthening in the γ′ strengthened nickel based superalloys.^11,13 In addition, coarse precipitations with black colour are distributed in the austenite. The EDS analysis (Fig. 1c) shows that the chemical composition of these precipitations is TiC (Table 2). The relatively strong pick of molybdenum in the EDS pattern reveals that Mo has participated with Ti to form (Ti,Mo)C combined carbide. These carbides have fall in the size range of 5–10 μm. These carbides can react with the austenitic binder to form low melting eutectic liquids, resulting in liquation cracking. Moreover, in some microstructural locations, a layered eutectic structures can be observed, which are attributed to γ–γ′ structures formed at the terminal stages of solidification process.^13,14 Based on the recent investigations, this eutectic structure can transform to liquid phase in the HAZ of the weld and causes the hot liquation cracks.^12,14,18

Table 2

Chemical composition of TiC precipitations

Element	Ti	Mo
Amount/wt-	58	42

Table 3

Chemical composition of eutectic structures enriched in Nb

Element	Nb	Ni	Cr	Fe	Other elements
Amount/wt-	19·5	50	15·7	14·5	Remained

Figure 2a indicates the microstructure of alloy 718 hot worked base metal. The presence of the mechanical twins within the microstructure clarifies which plastic deformation has been utilised for producing the alloy. The absence of the alloying elements segregation exhibits the success of the hot forging operations. The needle-like γ″-Ni₃Nb (or probably δ phase) precipitations have totally distributed within the austenite equiaxed grains. These precipitations have formed in the preferred crystallographic orientations intersecting each other at specific angles. The austenite grain boundaries are clearly thickened as a result of extensive accumulation of γ″. The studies have shown that the presence of this type of precipitations in the HAZ can intensify the susceptibility of the alloy to the liquation cracking.^10,13,28 In addition, coarse precipitations, maybe TiC, are dispersed within the austenite grains. Figure 2b illustrates the electron microphotograph of Alloy 718 base metal at higher magnification. The γ″ white coloured needle-like precipitations in the preferred orientations have appeared. These precipitations are the basic factors to strengthen the γ″ strengthened nickel based superalloys.^10,21

Microstructural investigation of weld metals

Figure 3a shows the microstructure of alloy 500 weld metal. The continuous network of dendritic boundaries is distinguished. The dendritic boundaries are brighter than the interior of the dendrites. Figure 3b has been taken for further identification of the Alloy 500 weld metal. The dendritic boundaries display the relatively significant segregation of the alloying elements. The EDS microanalyses for the interdendritic regions (Fig. 3c) and dendritic cores (Fig. 3d) reveal the high segregation tendency of titanium into the dendrite boundaries. Furthermore, elements such as Mo and Al, with lower tendency compared to Ti, have rejected the interdendritic liquids. Enrichment of the interdendritic regions in the segregating solutes can lead to remarkable reduction in the solidus temperature of the weld metal. It increases the brittleness temperature range, which causes the weld metal to be more prone to the solidification cracking. In the most interdendritic locations, microvoids that resulted from liquid to solid phase transformation shrinkages of the weld metal are seen. Coalescence of these microvoids due to the weld metal stresses can lead to propagation of hot solidification microfissures. Some hot microfissures are characterised in Fig. 3b.

Figure 4a is provided for identification of alloy 718 weld metal. Dendritic structure of the weld metal distinctly appeared with shiny dendrite boundaries. Figure 4b represents the higher magnification of the alloy 718 weld metal microstructure. The dendritic cores are darker and show lower amount of Nb (Fig. 4c). Two noteworthy microstructural features can be addressed in the dendrite boundaries: the segregation regions with white coloured appearance, which are enriched in Nb (Fig. 4d), and eutectic type morphologies, which display brighter colour. The brightness of these phases is attributed to the presence of high amount of Nb (Fig. 4e). Nb either in the interdendritic eutectic type structure or in the segregated state can greatly influence the weldability deterioration of the weld metal. The regions containing these Nb rich eutectic morphologies and segregated Nb in the austenite solid solution have lower solidification temperature, which increases the solidification temperature range of the weld metal.

Figure 5a indicates the microstructure of dissimilar weld metal with dilution level of 65. Although the microstructure is dendritic, the dendrite boundaries are not fully developed, as well as alloy 500 and alloy 718 weld metal microstructures. This can be explained by less tendency of alloying elements such as Nb and Ti to segregate to the final interdendritic liquids. The stronger segregation behaviour of alloying elements causes more developed dendritic structure. This structure is shown in higher magnification in Fig. 5b. In some locations, the segregation phenomenon is identified. Although the amount of Nb in the composition of the weld metal is significantly decreased, the formation of eutectic structures enriched in Nb has not been prevented. It seems that the less segregation and much more reduction in the formation of low melting eutectic structures are led to less susceptibility of the weld metal to solidification cracking, and as a result, the weldability has improved. In some eutectic constituent centres, shrinkage microvoids are seen, indicating that eutectics have been formed at the last stage of solidification process.

Dupont³⁹ and Moosavy et al.⁴⁰ could predict the susceptibility of nickel based superalloys to the hot cracking by utilising an analytical solidification model. This model has enabled comparing the superalloys together in order to rank them from the weldability viewpoint. The essential equation used to construct the model is Scheil equation in which the correlation between the k value of each element x composing the weld metal and the solid fraction f_S and the composition C has been considered (equation (1)) 1 By having the k values, the nominal composition of the elements and also some other microstructural, physical and operational parameters, the solidification range of the weld metal as an appropriate criterion to distinguish the hot weldability of superalloys studied here can be derived. The results of the calculations for the current work based on the model are presented in Table 4. These results demonstrate that the dissimilar weld with dilution level of 65 wt- has the minimum solidification temperature range and the least mushy zone extent, revealing the best weldability behaviour among the current investigated superalloys.

Table 4

Liquidus temperature, solidus temperature, solidification temperature range and extent of mushy zone for welds based on model calculations

Quantity weld type	Solidus temperature/°C	Liquidus temperature/°C	Solidification temperature range/°C	Extent of mushy zone/mm
Alloy 718	1090	1405	315	2·52
Alloy 500	1082	1400	318	2·54
Dissimilar weld	1181	1408	227	1·82

Microstructural characterisation of interfaces and HAZs

Figure 6a exhibits interesting features for the interface and HAZ of dissimilar weld between alloy 718 and alloy 500 with dilution level of 65, at the alloy 500 side. The alloy 500 base metal can be distinguished by coarse dendrites with dark regions of segregation at the top of the image. The bright dendritic regions, in the middle of the image, between the base metal and the fusion line, can be attributed to HAZ, in which relatively the main fraction of the area is free of segregation. High amount of heat received from the weld metal to the HAZ can lead to significant increase in the temperature. Therefore, the solubility limit and the diffusion rate of segregated elements in the austenite matrix were considerably enhanced. Consequently, the segregation elimination phenomenon occurs in the HAZ. Nevertheless, some of the interdendritic locations display strong segregation, which may have resulted from the formation of stable phases during the casting process. At these locations, especially those that are close to the interface between the weld metal and HAZ, the dendrite boundaries remelting and resolidifying occurred. These locations are suitable sites for nucleation and propagation of liquation cracks. Indeed, these liquated islands belong to an area that is named partially melted zone (PMZ). These resolidifed islands are the interdendritic regions of the base metal containing segregating elements, reacting precipitations and low melting eutectics, which have lower melting temperature in comparison to the surrounding areas. For more accurate identification of segregation phenomenon in the HAZ, Fig. 6b has been provided.

Figure 7 shows microstructural features for the interface and HAZ of dissimilar weld in the alloy 718 side. The HAZ clearly includes a region, in which the dissolution of γ″-Ni₃Nb or δ needle-like precipitations occurred (i.e. HAZ dissolution phenomenon). The on heating temperature cycle of HAZ causes these precipitations to decompose, and high cooling rate of HAZ during the on cooling cycle has prevented the further precipitation. According to the researches, redistribution of Nb in the HAZ austenite matrix can sensitise the weld to hot liquation cracking.^10,21,22 Figure 7b has shown the HAZ dissolution phenomenon by more detail. Although the heat received from the weld pool to the HAZ was able to eliminate the major segregation of alloying element in the interdendritic regions, the remained effects of the segregation can easily facilitate the propagation of liquation cracking.

Since the liquation in the HAZ of nickel based superalloys follows the mechanism of segregation and eutectic constituent melting,¹³ the suggested analytical solidification model mentioned in the section on ‘Microstructural investigation of weld metals’ can be utilised to analyse the hot liquation cracking in the HAZ of these alloys. According to the results obtained for the alloy weld metal solidification cracking in the section on ‘Microstructural investigation of weld metals’, it is expected that alloy 718 and alloy 500 exhibit relatively similar susceptibility behaviour to the liquation cracking. To discover a direct relation between the model findings on liquation cracking and the actual weldability behaviour of the current investigated superalloys, liquation cracking characterisation in the HAZ of the welds has been carried out. Figure 8a illustrates the HAZ and interface for the alloy 500 side, revealing that severe liquation cracking has occurred. It has passed through the weld metal, liquated islands of PMZ and then deeply propagated into the interior of the HAZ. Likely, the crack has nucleated at the liquated islands in the interdendritic region of PMZ and then propagates into the weld and base metals. The presence of segregated Ti in the interdendritic regions of the HAZ, and reacting precipitations and low melting eutectics enriched in Ti, in the PMZ islands, are the essential factors to induce this type of cracks in alloy 500. Figure 8b has displayed the hot liquation cracking in the HAZ of alloy 718 base metal. How the crack is elongated in the HAZ and propagated within the UZ, weld metal and base metal can be observed. Decomposed γ″-Ni₃Nb in the HAZ dissolution region, intergranular stable γ″-Ni₃Nb and segregated Nb to the boundaries play an important role to promote such hot liquation cracking. Furthermore, distinct view of unmixed zone appeared in the microstructure. This undesired region can be eliminated by utilising appropriate solutionising heat treatment.³⁷

Microhardness measurements

Figure 9 displays the results from the microhardness profile measurements in the dissimilar welds of alloy 718 to alloy 500. Based on the profile, the right and left sides are for alloy 500 and alloy 718 respectively. As observed, the microhardness average value for the weld metal is ∼300 HV. With coming to the alloy 500 interface, the hardness is increasing up to 360 HV. This increase in the hardness is attributed to the higher inherent strength and hardness of austenite solid solution of alloy 500 in comparison to alloy 718. By crossing from the weld interface and entering the HAZ, the hardness reduces and reaches the constant value of 300 HV in the alloy 500 base metal. The microhardness fluctuations in the alloy 500 base metal are related to the presence of the grain and dendritic boundaries and interdendritic intermetallic compounds such as γ′-Ni₃(Ti,Al), TiC and γ–γ′ eutectic structures, as seen in Fig. 1, that cause the local microhardness to increase. In the side of alloy 718, in the HAZ, the hardness has severely declined just after the weld interface and reached 220 HV. The significant decrease in the HAZ hardness can be interpreted with regard to Fig. 7. Indeed, the HAZ is the dissolution zone exhibited in Fig. 7, in which the γ″-Ni₃Nb participations disappeared. Since the γ″-Ni₃Nb particles are the chief reason of strengthening in alloy 718, dissolution of these compounds can lead to substantial decline of the hardness. Beyond the alloy 718 HAZ, the hardness is recovered again and reaches 360 HV due to secondary appearance of γ″-Ni₃Nb precipitations in the base metal. The microhardness values of the alloy 718 base metal are fluctuating between 360 and 380 HV that can be correlated with the non-uniform distribution of γ″-Ni₃Nb precipitations within the microstructure, as observed in Fig. 2.

Conclusions

For Alloy 500, the austenite matrix has been thoroughly occupied by γ′-Ni₃(Ti,Al). The needle-like γ″-Ni₃Nb precipitations have totally dispersed within the austenite equiaxed grains of alloy 718.

For the alloy 500 weld, titanium showed high segregation tendency. Coalescence of the microvoids led to hot solidification cracks. The eutectic type morphologies formed in alloy 718 weld are enriched in Nb. The microstructure of dissimilar weld metal displayed semideveloped coalescence of the dendrite boundaries.

The solidification model showed the best weldability for dissimilar weld with dilution of 65 compared to the alloy 500 and alloy 718 welds. The less segregation and significant reduction in the low melting eutectic structures caused improvement of the weldability in dissimilar weld.

The segregation elimination has occurred in the HAZ of alloy 500. In the PMZ, remelting and resolidifying occurred. These locations are sites for nucleation of liquation cracks.

For the alloy 718 HAZ, dissolution of needle-like γ″-Ni₃Nb occurred.

The HAZ and interface for the alloy 500 side revealed severe liquation cracking. The crack is initiated at the liquated islands in the PMZ. The hot liquation cracking in the HAZ of alloy 718 was observed.

The Vickers microhardness profile exhibited higher values for alloy 718 than alloy 500 base metal. Dissolution phenomenon in the HAZ of alloy 718 strongly decreased the hardness.

It is concluded that the control of dilution level can lead to crack free weld metals in the Inconel 718–Udimet 500 dissimilar joints. The present study offers the 65 dilution for sound weld metals. The dilution control is proposed as an effective approach for the fabrication of sound joints in such dissimilar welds.

Footnotes

Acknowledgement

The first author would like to thank the metallurgy team of Politecnico di Milano University, composed of Mrs S. Barella, Mr D. Mombelli and Mr A. Gruttadauria under supervision of Professor C. Mapelli, for their kind collaboration.

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