Elsevier

Journal of Manufacturing Processes

Volume 57, September 2020, Pages 36-47
Journal of Manufacturing Processes

Pre- and post-TLP bond solution treatments: Effects on the microstructure and mechanical properties of GTD-111 superalloy

https://doi.org/10.1016/j.jmapro.2020.06.011Get rights and content

Highlights

  • Standard heat treatment after bonding led to the uniformity of the composition and the distribution of γ'.

  • Applying the full solution annealing resulted in fully eliminating of the DAZ precipitation in the bond region.

  • After standard heat treatment in the non-isothermal solidification sample, MC and M6C carbides, η and eutectic phases were formed.

  • Standard heat treatment led to uniformity of hardness profile and increasing shear stress and toughness in the bonding district.

  • Mechanical properties had reduced in the non-isothermal solidification sample after heat treatment.

Abstract

This paper addresses the effects of bonding time and standard solution treatment on the microstructure and mechanical properties of GTD-111 superalloy bonded by transient liquid phase (TLP) using a Ni-7.58Cr-3.42Fe-3.6Si-2.95B amorphous interlayer. In this regard, three different heat treatment cycles namely HTsingle bondA, HTB and HTCsingle bondsingle bond were applied to separate TLP joints. In the HT-A cycle, full-solution annealing (1200 °C/4 h) of the base metal was first carried out followed by TLP bonding at 1120 °C/105 min, partial solution annealing (1120 °C/2 h) and finally aging treatment (845 °C/24 h). In the HTBsingle bond cycle, the first step was TLP bonding for 105 min at 1120 °C followed by full-solution, partial solution and aging heat treatments. Finally, the HTsingle bondC cycle was similar to the HTsingle bondB cycle except for the TLP bonding process which was performed at 1120 °C/15 min. The results indicated that HTsingle bondB treatment leads to a more uniform microstructure in the bond area and base metal compared to the microstructure obtained from HTsingle bondA and HTCsingle bond samples. Mechanical testing results indicated that a completely uniform hardness profile, similar to the base metal, was perceived in the bond region in the HTsingle bondB sample. The highest joint shear strength among all of the samples was obtained for the HTsingle bondB sample.

Introduction

GTD-111 is a cast nickel-based superalloy which has a multi-phase microstructure consisting of gamma matrix (γ), gamma prime precipitates (γʹ), carbides, gamma-gamma prime eutectic (γ/γʹ) and minor amounts of detrimental phases such as: δ, σ, ɳ and Laves [1]. Sajjadi et al. [2] have explained that there are two major strengthening mechanisms in GTD-111: 1) solid solution strengthening, 2) second-phase strengthening (i.e. precipitation hardening). The presence of elements such as W, Cr, Ti, Ta and Mo in the nickel matrix leads to solid solution strengthening. Sajjadi et al. [3] have reported that the formation of gamma prime precipitates with considerable amounts of Ni, Al, Mo and W in GTD-111 superalloy improves the strength of the alloy via precipitation hardening. The gamma prime precipitate is a superlattice phase with a general composition of Ni3(Al,Ti). There are two distinguishable structures of gamma prime phase in GTD-111 superalloy. The primary gamma prime with a cubic shape which is formed during solidification below 1200 °C and the secondary gamma prime with a spherical shape which nucleates and grows during the aging process after partial-solution treatment. Typical mechanical properties of GTD-111 include an ultimate tensile strength of 1011 MPa, hardness of 580 HV, shear strength of 750 MPa and high temperature strength of 854 MPa at 816 °C [4,5].

Shi et al. [6] have pointed out that GTD-111 is susceptible to erosion, corrosion, oxidation, thermal fatigue cracking and foreign object damage caused by drastic service conditions, which can lead to premature failure of the alloy, in spite of its remarkable properties. Huang and Miglietti [7] have reported that due to the high cost of turbine components manufactured with superalloys, it is economically favorable to repair the damaged components using techniques such as fusion welding, traditional brazing and diffusion brazing instead of manufacturing new components. Athiroj and Wangyao [8] have shown that fusion weld repaired GTD-111 and similar superalloy components are more vulnerable to mechanical property degradation due to heat affected zone (HAZ) cracking during the welding and post-weld heat treatments (PWHT). Moreover, in high temperature brazing, melting point depressants (MPDs) namely phosphorus, silicon and boron are added to the braze alloy in order to increase its fluidity. Philips et al. [9] have found that nevertheless the presence of these elements leads to the production of detrimental phases (e.g. borides, silicides and phosphides) during non-equilibrium solidification of the remaining liquid phase in the cooling stage. Adebajo [10] has claimed that the existence of intermetallic phases in the centerline of joints can have adverse effects on the overall performance of the component such as significant reduction in the mechanical properties, re-melting temperature and also corrosion and oxidation resistance.

Diffusion brazing (DB) also recognized as transient liquid phase (TLP) bonding is a superior technique in order to produce strong joints while overcoming the aforementioned limitations of prevalent joining methods [11,12]. Cook and Sorensen [13] have described that diffusion brazing process is carried out in four stages: (i) heating the bond setup to the bonding temperature and melting of the braze metal; (ii) dissolution of the substrate material after melting of the braze metal; (iii) isothermal solidification of the liquid phase in the joint when the liquid composition reaches the liquidus temperature; (iv) fulfillment of the isothermal solidification with diffusion of the MPD elements in to the substrate.

Heat treatment is carried out on bonded components for commercial applications to create a microstructure and elemental distribution in the bonding area similar to the base metal. The common heat treatment procedure for bonded superalloys consists of a heat treatment cycle for the TLP bond followed by a standard heat treatment for the base metal [[14], [15], [16]]. There is no doubt that it is economically preferable to be able to simultaneously carry out the TLP bonding heat treatment and standard heat treatment cycles to save time and resources.

There are many reports on the effects of joining parameters on the microstructure and mechanical properties of the bond. Arhami et al. [17] have reported that increasing bonding time results in the growth of a solid solution phase (e.g. γ in superalloys) and also avoids the formation of adverse brittle intermetallic phases in the bond area in a superalloy system. Baharzadeh et al. [18] have investigated dissimilar joining of IN X-750 to SAF 2205 by TLP process with BNi-2 interlayer foil. They have reported that raising of bonding temperature will increase the isothermal solidification rate and reduce the bonding time. According to the results of Hadibeyk et al. [19], the width of the diffusion affected zone (DAZ) and the athermally solidified zone (ASZ) are raised by increasing the thickness of the filler metal. However, little research can be found on the application of standard heat treatment cycles on the TLP bonded joints, or on the combination of TLP bonding and standard heat treatment. Pouranvari et al. [20] have reported that using the standard heat treatment on the TLP joint can eliminate the boride precipitates in the DAZ and increase the amount of dissolved Nb + Al + Ti in the isothermally solidified zone (ISZ) leading to the improvement of overall mechanical properties. Shakerin et al. [21] have illustrated that applying standard heat treatment on TLP bonded IN-738 leads to increased Ti and Al content, as the base alloying elements, in the ISZ region and enhances the overall mechanical properties. However, there is no research about the effect of full solution annealing after diffusion brazing on the microstructure and mechanical properties of the bond.

In this paper, the effects of standard solution annealing (partial and full-solution) and aging heat treatment of GTD-111 superalloy, before and after TLP bonding, on the microstructure and mechanical properties are investigated. In addition, the effect of a simultaneous post-TLP bonding and base metal heat treatment process on the microstructure and mechanical properties is also studied for the first time.

Section snippets

Experimental procedure

As-cast GTD-111 was used as the base metal in this research. Amorphous Ni-Si-B-Fe-Cr interlayer (AWS BNi-2) with a thickness of 25 μm was applied as the filler metal. The chemical composition of the substrate and filler material are listed in Table 1. Specimens with dimensions of 10 × 10 × 5 mm were sectioned from the base metal using electro-discharge machining (EDM). All surfaces of the test coupons were polished by means of silicon carbide grinding papers (from 100 to 800 grit) and

Microstructure of base metal before and after standard heat treatment

Fig. 3 shows the most common phases present in the microstructure of GTD-111 superalloy in the as-cast or pre-standard heat treatment condition. As shown in Fig. 3, the GTD-111 includes different phases such as γ matrix, γˊ precipitates, γ - γˊ eutectic islands, carbides and topologically closed packed (TCP) phases. The mole fraction of equilibrium phases formed in the base metal were calculated using Thermo-Calc Software TTNI7 database and are presented in Fig. 4. According to this figure,

Conclusions

TLP bonding of GTD-111 was carried out at 1120 °C for two different times, according to incomplete and complete isothermal solidification. Standard solution annealing was performed as pre- or post-TLP bond treatment. The microstructural and mechanical study illustrated the following results:

  • 1

    γ' precipitates were observed before and after solution annealing and also after aging in the form of non-geometric, spherical, blocky structure, respectively. Very small and spherical secondary γˊ particles

Acknowledgements

The authors appreciate the support from Ferdowsi University of Mashhad (FUM) under the research scheme No. 2/43941. The authors also acknowledge the technical assistance on Thermo-Calc Software calculations from the Canadian Centre for Welding and Joining (CCWJ) at University of Alberta.

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