Microstructural response to heat affected zone cracking of prewelding heat-treated Inconel 939 superalloy

M A TE RI A L S C H A RAC TE RI ZA T ION 6 2 ( 20 1 1 ) 1 1 1 6–1 1 2 3

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Microstructural response to heat affected zone cracking of prewelding heat-treated Inconel 939 superalloy M.A. González a,⁎, D.I. Martínez a , A. Pérez a , H. Guajardo b , A. Garza c a Facultad de Ingeniería Mecánica y Eléctrica (FIME-UANL), Av. Universidad s/n. Ciudad Universitaria, C.P.66451 San Nicolás de los Garza, N.L., Mexico b FRISA Aerospace, S.A. de C.V., Valentín G. Rivero No. 200, Col. Los Treviño, C.P. 66150, Santa Caterina N.L., Mexico c Corporación Mexicana de Investigación en Materiales S.A. de C.V. (COMIMSA), Ciencia y Tecnología No.790, Saltillo 400, C.P. 25295 Saltillo Coah., Mexico

AR TIC LE D ATA

ABSTR ACT

Article history:

The microstructural response to cracking in the heat-affected zone (HAZ) of a nickel-based

Received 26 April 2011

IN 939 superalloy after prewelding heat treatments (PWHT) was investigated. The PWHT

Received in revised form

specimens showed two different microstructures: 1) spherical ordered γ′ precipitates

19 September 2011

(357–442 nm), with blocky MC and discreet M23C6 carbides dispersed within the coarse den-

Accepted 21 September 2011

drites and in the interdendritic regions; and 2) ordered γ′ precipitates in “ogdoadically” diced cube shapes and coarse MC carbides within the dendrites and in the interdendritic re-

Keywords:

gions. After being tungsten inert gas welded (TIG) applying low heat input, welding speed

Welding

and using a more ductile filler alloy, specimens with microstructures consisting of spherical

Ni-based superalloy

γ′ precipitate particles and dispersed discreet MC carbides along the grain boundaries, dis-

Microstructure

played a considerably improved weldability due to a strong reduction of the intergranular

SEM

HAZ cracking associated with the liquation microfissuring phenomena.

Grain boundary

© 2011 Elsevier Inc. All rights reserved.

HAZ microfissuring

1.

Introduction

Continuous attempts to increase the overall efficiency and performance of modern power generation systems by operation at higher temperatures has demanded significant progress in the development of superalloys that combine corrosion resistance with long-term mechanical strength. Cast high-chromium, nickel-based superalloys such as IN 939 are of particular importance and are widely used in the blades and vanes of land-based and marine gas turbines for their remarkable superior hot corrosion and oxidation resistance due to the higher chromium content than the IN 738 alloy. This superalloy contains a considerable volume fraction (~60%) of ordered L12 γ′ intermetallic compounds of the type (Ni3,Co3) (Al, Ti or Ta), which is

the main high-temperature strengthening phase [1–3]. The blades of these types of turbines frequently suffer from premature failures such as airfoil tip loss due to hot corrosion and low-cycle fatigue cracking. The high replacement cost of these parts creates a demand for the refurbishment of damaged turbine airfoils at a lower cost. The fusion welding process is considered one of the most desirable and economical processes for commercial repair applications [4–5]. However, it is well known that precipitationhardened nickel-based superalloys such as IN 939 possess poor weldability, because these alloys are highly susceptible to microfissuring during the welding cycle, predominantly in the heat-affected zone (HAZ) [4,6] due to their considerable content of gamma prime (γ′) formers such as Al and Ti (≥6 wt.%).

⁎ Corresponding author. Tel.: + 52 844 411 3200; fax: +52 844 415 2151. E-mail addresses: [email protected] (M.A. González), [email protected] (D.I. Martínez), [email protected] (A. Pérez), [email protected] (H. Guajardo), [email protected] (A. Garza). 1044-5803/$ – see front matter © 2011 Elsevier Inc. All rights reserved. doi:10.1016/j.matchar.2011.09.006

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Several studies have been conducted with similar superalloys in attempts to eliminate or minimize the HAZ cracking of welds. Shaw et al. [7] described a method for considerably improving the weldability of this alloy through prewelding heat treatment (PWHT); however, there is limited information available related to the mechanism of HAZ cracking in this alloy. In contrast, there is copious information available for other alloys from studies related to their cracking behavior, reporting that intergranular cracking is affected by the following: (1) the presence of precipitates on grain boundaries (e.g., carbides, borides, gamma prime and other low-melting phases) that can liquate and initiate cracking; and (2) the morphology of grain boundaries (planar serrated), which can affect the ability of an intergranular crack to propagate [6–9]. HAZ liquation is known to occur either by nonequilibrium interface melting below the alloy's solidus or by equilibrium supersolidus melting [8,10]. Several authors have reported improvements in the HAZcracking resistance of many superalloys with the use of a special TIG welding process preceded by a preheating treatment [6,8,10]. This behavior is attributed to the creation of a microstructure that is capable of relaxing the stresses generated during the welding process, considerably reducing intergranular liquation microfissuring. Therefore, understanding the evolution of microstructures in the HAZ in response to the rapid heat pulses of welding and the relation between preand postwelding microstructures is imperative. The aim of the study was to understand the effect(s) of prewelding heat treatments on the metallurgical reactions that contribute to HAZ-liquation cracking during the welding of IN 939 superalloy. Herein, we present the results of this investigation and discuss their relation to HAZ microfissuring.

2.

Experimental Procedures

The materials used in the present study were cast turbine blades made of Inconel 939 superalloy. The chemical composition was (wt.%) 0.14 C, 23.23 Cr, 19.31 Co, 1.83 W, 0.16 Mo, 1.04 Nb, 0.20 Fe, 1.85 Al, 3.83 Ti, 1.37 Ta, 0.10 Zr, and the balance nickel. Test coupons (30 × 5 mm) were cut and machined from the airfoils of the blades. The specimens were subjected to several prewelding heat treatments (Table 1) including solution treatment (1160 °C/4 h/FC), as suggested by Shaw et al. [7], and a proposed alternative solution treatment (1145 °C/4 h/FC (furnace cool), designated AST). Both samples were furnace-cooled at a rate of 6 °C/min until they reached the subsequent aging isothermal; this last treatment was held for 16 h and then FC to room temperature. The heat-treated specimens were surface ground and cleaned with acetone to remove surface contamination and then tungsten inert gas (TIG) arc welded, in a single pass using A230 filler weld (Table 2) which is a standard low-strength (nonprecipitation hardening) welding at a constant welding speed of 45 mm/min and a current (I) of 25 A DC at 10 V using argon as a shielding gas. The heat input was 0.33 kJ/mm, as calculated using the expression HI = (V ×I × 60 /S ×1000), where S is the welding speed. Previous researchers have shown that low- and intermediate-strength fillers reduce microfissuring [8,11,12]. Welded specimens were sectioned transversely to the welding direction with a diamond cutting wheel to produce five sections from each test coupon, which were subsequently prepared

Table 1 – List of prewelding heat treatments. Identification H2 H3 H4 H5 H6 H7 H8 H9 FC-furnace cooled

Heat treatment 1160 °C/4 h/FC + 1000 °C/16 h/FC 1160 °C/4 h/FC + 950 °C/16 h/FC 1160 °C/4 h/FC + 900 °C/16 h/FC 1160 °C/4 h/FC + 850 °C/16 h/FC 1145 °C/4 h/FC + 1000 °C/16 h/FC 1145 °C/4 h/FC + 950 °C/16 h/FC 1145 °C/4 h/FC + 900 °C/16 h/FC 1145 °C/4 h/FC + 850 °C/16 h/FC

by standard metallographic procedures for microstructural examination. The microstructures of the as-received material, the prewelding heat-treated alloy and the weld HAZ were evaluated by optical and scanning electron microscopy (SEM) using a JEOL-6490 equipped with an Oxford energy dispersive X-ray spectrometer with an ultrathin window detector. For this purpose, the specimens were electrolytically etched in a solution of 10 mL H3PO4 +50 mL H2SO4 +40 mL HNO3. The base-alloy, HAZ and fusion-zone (FZ) hardness were determined by a Vickers hardness tester under a load of 0.5 kg.

3.

Results and Discussion

3.1. Microstructure of As-received Base Alloy, Prewelding Heat-treated and Welded Specimens The received material had been operating for almost 40,000 h at temperatures near 850 °C in the hot section of a gas turbine. The microstructure of the blades' airfoil after this period of operation consisted of elongated serrated grains, approximately 205 μm in diameter, contained in a cored dendritic structure, indicating that the blades were conventionally cast. Second reaction precipitates (MC carbides) showed script-type and coarse blocky morphologies and were distributed mainly in the interdendritic regions; some of these carbides were preceded by small continuous films of M23C6 carbides and large (~1.5 μm) primary gamma prime (γ′) overage particles along the grain boundaries (Fig. 1), which indicates that the following reaction took place: MC+ γ → M23C6 + γ′ [13]. Fig. 2 shows that, as a result of the exposure of the alloy to high temperature for long Table 2 – Composition of weld filler alloy, wt.%. Element

Alloy A230

Ni C Al Si Ti Cr Mn Fe Co W Mo Nb Zn

Bal. 0.05 0.41 0.24 0.42 21.8 0.50 0.68 0.017 13.5 0.98 0.002 0.37

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Fig. 1 – SEM micrograph of the as-received alloy.

times, coarsening and agglomeration of the primary γ′ precipitate particles clearly increased at expense of the small spherical particles known as secondary γ′. The clear disparity in particle size between the dendritic and interdendritic regions and the presence of several nonequilibrium γ–γ′ eutectics (Fig. 1) were attributed to the diffusion of Al and Ti (γ′ formers) to the dendritic core [10,14,15]. The heat treatments H2, H3, H6 and H7 showed a complete dissolution of all this intermetallics indicated earlier, resulting in a bimodal microstructure of homogeneously distributed primary γ′ spheroids (~0.35 μm) and smaller spheroids of secondary γ′ (Fig. 3a) in both the dendritic core and the interdendritic regions. In addition, several secondary solidification constituents such as MC carbides were observed to be mostly of the blocky type and were distributed in both the dendritic core and along the grain boundaries. Particles of M23C6 carbides were also observed (Fig. 3b). All this suggests the dissolution and the reprecipitation of the initial core and script-type carbides during the solution treatment [1,7,15]. Chemical analysis of different MC carbides with SEM-EDS (Figure 4) showed that Ti, Nb and Ta are the principal compound elements of these type of carbides, other elements such as Cr and Ni also appear in the spectrums because even if the analysis was done at specific microareas of the carbide particles, some of the adjacent material is

Fig. 2 – Micrographs of primary and secondary γ′ particles in the as-received alloy.

Fig. 3 – SEM micrographs of the alloy after PWHT: (a) distribution of the primary and secondary γ′; (b) MC carbide distribution. also analyzed. The nonequilibrium γ–γ′ eutectics were observed to dissolve completely, and the formation of detrimental particles such as the η phase were avoided with these treatments, in accordance with the results obtained by Shaw et al. [7]. A complete dissolution of the primary and secondary γ′ precipitates was also observed with heat treatments H4, H5, H8 and H9. The

Fig. 4 – EDS of the MC-type carbides.

M A TE R IA L S C H A RAC TE RI ZA T ION 6 2 ( 2 01 1 ) 1 1 1 6–1 1 2 3

lower aging temperatures of these treatments resulted in microstructures of “ogdoadically” diced cube-shaped γ′ particles (Fig. 5a), which represent clusters of cubic particles seen as four planes in the {001} directions [16]. Isolated coarse carbides and nonequilibrium γ–γ′ eutectics were observed to persist along the grain boundaries (Fig. 5b) with treatments H8 and H9. This is attributed to the lower solution temperature used in these treatments. An overall SEM micrograph of the cross-section of the TIGwelded specimens is shown in Fig. 6. A fine interdendritic microstructure is observed in the fusion zone (FZ). Treatments H2, H3, H6 and H7, exhibited some liquid film migration (LFM) features (Fig. 7), limited to the immediate vicinity of the fusion boundary. Apparently uncracked liquated grain boundaries were possible due to either insufficient welding stresses across such boundaries during weld cooling or to the healing of the resulting cracks due to backfilling from a relatively large volume of intergranular liquid. However, it should be mentioned that although the liquated grain boundaries such as the one shown in Fig. 7 apparently remained uncracked during welding, such heterogeneous regions can exhibit nonuniform deformations during the stress relaxation of a postwelding heat treatment. This might result in crack

Fig. 5 – SEM micrographs of the alloy after PWHT (aging temperatures 900 °C and 850 °C): (a) distribution of γ′ particles; (b) incomplete dissolution of grain-boundary MC carbides and γ–γ′ eutectics.

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Fig. 6 – SEM micrograph of the cross-section of the single-pass TIG-welded specimens (backscattered electron image).

initiation and/or propagation along such grain boundaries due to excessive local strain concentration. It has been demonstrated that these liquated regions in Ni-based superalloys are more susceptible to strain age cracking [15–18]; however, this study focused only on explaining cracking developed during the welding. Microcracking was observed in the HAZ of welded specimens with the H4, H5, H8 and H9 treatments; the cracks originated slightly away from the fusion boundary and in some cases, they extend as much as 220 μm into the base alloy and occasionally partly into the FZ (Fig. 8a). The cracks were intergranular and displayed a relatively irregular and jagged path showing the typical character of liquation cracks. The resolidified constituent surrounding the γ′ precipitate particles (Fig. 8b) along the grain boundary interface was observed to penetrate into the grain boundary, promoting the microfissuring of these regions. Some of the microfissures were observed to appear at liquated grain boundaries that exhibited LFM features (Fig. 9), which suggests the formation of a grain boundary liquid during the welding thermal cycle. The liquid was observed to be present in the form of thick, continuous layers, particularly along the grain-boundary regions in the immediate vicinity of the FZ. The microfissuring in these materials is generally attributed to the inability of

Fig. 7 – SEM micrograph showing LFM near the FZ.

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the liquated intergranular regions to accommodate tensile stress developed during the on-cooling cycle, so the available strength is exceeded, resulting in decohesion across one of the solid–liquid interfaces [10,14].

3.2. Effect of Intergranular Constitutional Liquation on HAZ Microfissuring An extensive grain-boundary liquation was observed in addition to the supersolidus liquation that is expected to occur in all weldments due to heating above the equilibrium solidus temperature. This phenomenon was first proposed by Pepe and Savage [19] and has been subsequently observed by other investigators in various alloy systems involving the main strengthening phase (γ′ precipitates) of nickel-based superalloys hardened by precipitation as well as in carbides, borides and sulfide particles [6,14–18]. It has been attributed mainly to subsolidus liquation of second phases present in the prewelded alloy [6,18] due to a eutectic-type liquation reaction between the second-phase particle and the matrix, producing a nonequilibrium solute film (rich in gamma formers) at the interface. The basic requirement for the occurrence of constitutional liquation of second-phase intermetallic compounds AxBy is rapid heating of the alloy during welding to

Fig. 8 – SEM micrographs of: (a) Jagged path typical of HAZ liquation cracks; (b) resolidified products associated with intergranularly liquated MC carbides and γ′ particles.

Fig. 9 – SEM micrograph of microfissuring at LFM and liquated HAZ grain boundaries associated with resolidified products.

the eutectic temperature or higher. As a result, the precipitate does not have sufficient time to dissolve completely in the matrix [20–21]. Therefore, the susceptibility of second phases to constitutional liquation can be related to its solid-state dissolution behavior and its complete dissolution prior to reaching the eutectic temperature precludes the occurrence of liquation. Previous researchers in similar alloy systems have reported how the initial particle size, morphology and heating rate affect the integrated time required for homogenization by the diffusion process during continuous heating and how this can cause γ′ precipitates in nickel-based superalloys to persist at temperatures above their equilibrium solvus temperature [14,21–22], they found that if the initial particle presents a ogdoadically-diced cube shapes (irregular shape) and coarse size, the temperature to reach the complete dissolution tend to increase. In the present work, it was observed that in the rapid heating typical of the TIG welding process, the dissolution behavior of γ′ particles significantly deviated from equilibrium. This occurs because the particles survive beyond their equilibrium dissolution temperature to temperatures at which they can constitutionally liquate with the surrounding γ matrix by a eutectic-type reaction (Fig. 10). As mentioned previously, at a specific heating rate, particle size and morphology directly affect the tendency of γ′ particles to survive during continuous heating to temperatures at which they can constitutionally liquate and contribute to grain-boundary liquation. It was observed that mainly the specimens that presented “ogdoadically” diced cubic γ′ particles showed both intergranular and intragranular constitutional liquation, this behavior as mentioned earlier is closely related to the irregular shape and the initial particle size of the precipitates, increasing the tendency to persist and liquate at temperatures above their equilibrium solvus temperature (Fig. 10). In contrast, the specimens with spherical particles showed no evidence of constitutional liquation. It has been previously demonstrated that for a planar precipitate, the time required for complete homogenization is considerably larger (~two orders of magnitude) than the time required for precipitate dissolution. For a spherical geometry, the concentration in the center of the dissolved precipitate at the instant of its dissolution decreases to a value lower than the equilibrium solubility value at the matrix–precipitate interface [23] increasing

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Fig. 10 – SEM micrograph of the HAZ microstructure showing constitutional liquation of γ′ particles.

the dissolution rate of the precipitate; as a result during fast heating the spherical precipitates will dissolve almost completely before reaching their equilibrium solvus temperature. Careful examination of the HAZ cracks showed that besides the liquation of γ′ particles, MC carbides (Fig. 11a and b), which are secondary solidification-reaction products, also presented liquation. As mentioned earlier, weld HAZ cracking results from a competition between the mechanical driving force for cracking due to tensile welding stresses and metallurgical factors such as intergranular liquation, which causes the embrittlement of grain boundaries. Furthermore, as can be seen from the expression σ = 2γSL/h, first suggested by Miller et al. [24] (where σ is the tensile stress required to overcome the surface tension γSL on a boundary containing a liquid film of thickness h), any parameter that decreases the thickness of the intergranular liquid film would enhance resistance to HAZ cracking by increasing the stress required to cause liquation microfissuring. As stated earlier, in the present work, the prewelding heat-treated specimens which showed less and severe grain-boundary liquation; in both cases the extension and thickness of the liquid films were similar and differed from the expected behavior in line with Miller's equation, this behavior can be attributed to the low volume fraction of liquating γ′ (~0.4 μm) precipitates. Comparing with the considerable volume fraction of liquating γ′ (~0.8 μm) precipitates observed in similar alloys [14,17] which have shown several intergranular liquation cracking. For alloys similar to IN 939, such as IN 738, other investigators have reported that with different prewelding heattreatment conditions, welding stresses of similar magnitude occur during cooling in the HAZ [14–15,22]. The magnitude of welding stresses in the HAZ, however, can be significantly influenced by the ability of the alloy to relax part of the stress, which is related to the hardness of the material prior to welding. It has been demonstrated for similar alloys such as IN 738 [22,25] that a hard base alloy showing poor stress relaxation can cause large stresses to be concentrated in the relatively weakened HAZ. Conversely, a base alloy with lower hardness that is capable of substantial stress relaxation can reduce the stresses in the HAZ, in amounts sufficient to reduce the driving force for

Fig. 11 – SEM micrographs of: (a and b) HAZ cracking at constitutionally liquated MC carbides.

intergranular liquation cracking. Fig. 12 shows the hardness variation for the alloy IN 939 with prewelding heat treatments under various conditions. As shown here, the hardness was very similar for all treatments; these results indicate that heat-treated specimens that presented only limited liquation near the FZ were able to relax the stresses in the HAZ, resulting in seemingly uncracked grain boundaries. Therefore, in the present study, we observed the prevention or diminishing of HAZ microfissuring as a result of grain-boundary liquation in samples showing spherical γ′ precipitate particles, and no more than three cracks were observed in the specimens with severe liquation. The low quantity of cracks in these specimens was attributed, as mentioned earlier, to the prewelding hardness of the alloy being capable of accommodating the tensile stresses generated in the HAZ during the welding thermal cycle.

3.3.

Influence of Welding Parameters on HAZ Cracking

The amount of stress generated during the welding process can be controlled to a certain extent by selecting the proper welding parameters. It has been reported that a lower welding speed, a low heat input and the use of a more ductile filler alloy can significantly affect the amount of stress generated during fusion-repair processes [4,8,11,22,26–27]. As indicated

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ductile filler alloy. The interaction of these factors led to a lower residual stress concentration and directly influenced the capability of the weld to accommodate the stresses generated during the on-cooling cycle. Therefore, the study of precipitation-hardening filler alloys using similar parameters of heat input and welding speeds is equally relevant.

4.

Fig. 12 – Hardness of base alloy in the prewelding heat-treated conditions. previously, in our experimental procedure, the selected parameters such as speed (45 mm/min) and heat input (0.33 kJ/mm) are considered low in comparison with similar investigations on alloys such as IN-738 [4,11,27]. Both factors influence the thermal gradient; the welding speed determines the fraction of time (t) in which the TIG torch spends on depositing energy into a section of the workpiece; as welding speed is increased, t is reduced, and less energy is deposited into the workpiece. The heat energy deposited in the material not only heats the area directly under the weld, but some of the heat input energy is conducted into the adjacent portion of the base metal. The more time the torch spends in each section, the more heat is conducted into the base metal, resulting in a lower thermal gradient. The magnitude of welding stress is known to be directly proportional to the HAZ thermal gradient. The results from the present study show that besides the influence of the prewelding microstructure (hardness) in the appearance of HAZ microfissuring, the choice of welding parameters can create a lower level of residual stresses and consequently reduce HAZ cracking tendencies. A considerable softening of the HAZ has been observed for lower heat inputs [4]. Fig. 13 shows the hardness distribution between the base metal and the filler alloy; a softening of the HAZ was clearly observed, which can also contribute to the HAZ microfissuring tendency. The low percentage of cracking observed in the specimens with microstructures of γ′ spheroids and discreet blocky MC carbides can be strongly influenced by the low heat input and welding-speed parameters as well as the use of a more

Summary and Conclusions

1. The prewelding heat treatments resulted in the following: 1) a homogeneous microstructure of bimodally distributed γ′ spheroids and smaller secondary γ′ spheroids and secondary solidification with block-type MC carbides; and 2) a homogeneous microstructure of “ogdoadically” diced cubic γ′ particles and isolated coarse MC carbides. 2. The complete dissolution of the primary and secondary γ′ precipitate particles was observed with the alternative solution treatment (ATS); however, isolated γ–γ′ eutectics remained undissolved and can contribute to HAZ liquation cracking during welding. 3. Intergranular HAZ liquation cracking of the TIG welded IN 939 superalloy was generated by liquation reactions of secondary microconstituents (MC) and γ′ precipitate particles. 4. The prewelding specimens that produced less HAZ grainboundary liquation cracking were those that showed γ′ spheroids and discreet blocky MC carbides, compared with those that exhibited an “ogdoadically” diced cube γ′ particles and coarse MC carbides, which exhibited a higher degree of both intergranular and transgranular liquation. 5. Also the low susceptibility to grain-boundary liquation cracking of prewelding microstructures (γ′ spheroids and discreet blocky MC carbides) can be attributed to the hardness of the base alloy, which was capable of relaxing the stress generated during weld cooling. 6. This behavior may also be related to the low heat input, welding speed and the use of a more ductile filler alloy, which together influenced the capability of IN 939 welds to accommodate the stresses generated during the oncooling cycle, reducing HAZ cracking susceptibility.

Acknowledgments The authors would like to thank the CONACyT Mexico and the Paicyt as well as the Universidad Autonoma de Nuevo León for their financial support.

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Fig. 13 – Hardness distribution in TIG welds.

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