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Microstructure and mechanical properties of transient liquid phase (TLP)-bonded Ni3Al intermetallic compounds
Materials Today Communications 21 (2019) 100619
Contents lists available at ScienceDirect
Materials Today Communications journal homepage: www.elsevier.com/locate/mtcomm
Microstructure and mechanical properties of transient liquid phase (TLP)bonded Ni3Al intermetallic compounds
T
M. Soltani Samani, A. Bahrami , F. Karimzadeh ⁎
Department of Materials Engineering, Isfahan University of Technology, Isfahan 84156-83111, Iran
ARTICLE INFO
ABSTRACT
Keywords: Transient liquid phase bonding (TLP) Ni3Al Joining Intermetallic AWS BNi-2 Microstructure
This paper investigates microstructure and mechanical properties of transient-liquid-phase (TLP) bonding for the joining of Ni3Al intermetallic compound. The jointing was carried by using AWS BNi-2 interlayer in a vacuum furnace at 1050 °C. Effects of holding times for 30, 60, 90, and 120 min on the microstructure and mechanical properties of the joint have been investigated. The microstructure of samples was evaluated using optical and scanning electron microscopes (SEM). Phase identification was carried using energy dispersive spectroscopy (EDS) and X-ray diffractometer (XRD). Mechanical properties of joints were studied by hardness profilometry across the interface as well as shear test. Results show that no isothermally solidified zone has been formed in holding time up to 120 min. Also, solidification cracking was observed in the centerline of joints. EDS analysis showed that the centerline is rich in boron and some boride intermetallic compounds. Increasing holding time is associated with a decrease in the amount of brittle phases in the athermally solidified zone. Microhardness profilometry across the interface show that the maximum hardness is attained in the middle of the joint. This is attributable to the fact that the center of the joint is rich in hard nickel and chromium boride phases. Results also show that both hardness and shear strength of the joint increase with increasing holding time. Fractography results indicate that fracture takes place in a brittle mode.
1. Introduction Intermetallic compounds are known as materials with unique characteristics, including high temperature strength, high temperature oxidation/corrosion resistance, and excellent structural stability at high temperatures. This makes these alloys excellent candidates for many demanding applications in advanced industries. Amongst these intermetallic compounds, Ni3Al intermetallic compound is one of the most widely used ones, owing to the fact that it has a comparatively better formability. Ni3Al intermetallic compound is currently used as a commercial alloy in compressors and turbine blades in different industries [1,2]. A major obstacle in further using of this intermetallic alloy is the welding/joining of this compound. This difficulty originates from the inherent brittleness of intermetallic compounds. Obviously, amongst different fusion and solid state-based welding techniques [3–5], those which induce lower residual stress and heat input in the weldment are more preferred. Fusion-based welding processes, including gas tungsten arc welding (GTAW), electron beam welding (EBW) and laser beam welding are intrinsically prone to solidification-induced defects (cavities, solidification cracking, incomplete welds, voids, …). Given that intermetallic alloys are mostly used in very sensitive and demanding ⁎
applications, hardly is there any tolerance for mentioned defects, inferring that fusion-based techniques must be used with extreme precautions. In an attempt to minimize solidification cracking in laserwelded Ni3Al alloy, Zhang and et al. [6] postulated that controlling the heat input is a key controlling issue, when it comes to the sensitivity of cracks in the heat affect zone (HAZ) area. Obviously, the coarse-grain structure is more sensitive to solidification cracking during welding. Santla and et al. [7] did the same study during the GTAW process. They postulated that controlling the chemistry of the alloy is very influential in minimizing solidification cracking. Contrary to fusion welding techniques, solid-state-based methods, like friction welding process, have much less heat inputs. On the other hand, the downside of friction welding is that the workpiece undergoes excessive plastic deformation, making this method less applicable for the obvious reason of brittleness of intermetallic compounds. With that said, transient liquid-phase process (TLP) appears to be very promising. TLP, known to be a combination of diffusion bonding and brazing, is based on the formation of a transitional melting phase at the interface. Diffusion of alloying elements across the interface during TLP, while temperature is kept constant, ultimately results in the formation of the joint [8]. Microstructure and mechanical properties of TLP joints largely depend on the type and
Corresponding author. E-mail address: [email protected] (A. Bahrami).
https://doi.org/10.1016/j.mtcomm.2019.100619 Received 28 May 2019; Received in revised form 25 August 2019; Accepted 26 August 2019 Available online 27 August 2019 2352-4928/ © 2019 Elsevier Ltd. All rights reserved.
Materials Today Communications 21 (2019) 100619
M. Soltani Samani, et al.
Fig. 1. Schematics of the holder for the shear test.
Fig. 2. Dendritic structure of as-cast Ni3Al compound.
the chemical composition of the interlayer. Normally, melting point depressant elements (e.g. boron) are present in the interlayer to reduce the melting point of the interlayer. The lower the melting point, the easier is the formation of isothermal solidification zone at the joint [9]. Ojo and et al. [10] joined Ni3Al-based single crystal intermetallic alloy using TLP. They showed that insufficient holding time results in the appearance of undesirable brittle eutectic compounds at the joint centerline. They postulated that a mixture of brazing composite powder and interlayer can be used to reduce the time to complete isothermal solidification. However, their findings are limited to single crystals. To our knowledge, hardly is there any comprehensive study on the microstructure and mechanical properties of TLP-joined polycrystalline Ni3Al alloys. This paper investigates the TLP joining of polycrystalline Ni3Al intermetallic compound, using AWS BNi-2 as interlayer. AWS BNi-2 contains B and Si, with both are known to be melting point depressant elements. Besides, AWS BNi-2 has reportedly a very good corrosion resistance and creep strength [11]. In this study, the effects of holding time at a constant temperature/pressure on the mechanical properties and microstructure of bonding are also investigated. Hardly is there any comprehensive report on the microstructure and mechanical properties of TLP-bonded Ni3Al intermetallic compounds in the literature. This paper is an attempt to investigate how properties of TLPbonded Ni3Al intermetallic compounds can be controlled by changing the microstructure. 2. Experimental procedure
Fig. 3. XRD spectrum of as-cast Ni3Al compound, in which peaks coincide with Ni3Al peaks.
The Ni3Al intermetallic compound was first synthesized, in a
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Fig. 4. Effects of holding time for a) 30, b) 60, c) 90, and d) 120 min on the microstructure of the joint.
vacuum arc remelting (VAR) furnace, using pure nickel (99%) and aluminum rod (99.9%). The VAR was carried out three times on each specimen to make sure that alloying elements were homogeneously mixed. Experimental samples for joining were cut in 3 mm thick 10 × 15 mm pieces using electro discharge machining (EDM). TLP was performed in YMVR 1700 vacuum furnace under vacuum of 10−5 bar. The 50 μm thick interlayer AWS BNi-2, used in this study, had the following composition: Ni-5.9Cr-7.0Si-2.8Fe-3B (wt.%). The TLP was performed at 1050 °C four different holding times: 30, 60, 90, and 120 min. Samples were first ground before joining with polishing paper 1200. Both the interlayer and samples were sonicated in acetone for 25 min to ensure that surfaces were not contaminated. A fixture with dimensions of 10 × 20 × 70 mm was used to hold samples together during TLP. Microstructural studies were carried out using optical and electron microscopies. Samples were first ground with grinding papers, followed by polishing with SiC polishing paste. Marble solution (0.5 HCl+H2O+0.5 CuSO4) was used to etch samples. Energy dispersive XRay spectroscopy (EDS) was used to identify phases in the microstructure and to do the line scan across the interface. Vickers hardness profilometry was conducted across the interface, using 100 g load. To
determine the shear strength of the joint, samples of 10 × 10 mm were cut and placed in a fixture, as schematically shown in Fig. 1, then they were then loaded until they broke. The loading speed was chosen to be 1 mm/min. Finally, scanning electron microscope (SEM) was used to do the fractography on the fracture surface (Fig. 2). 3. Results and discussion 3.1. Microstructure of base metal Fig. 3 shows the microstructure of as-cast Ni3Al compound, in which a dendritic structure is clearly visible. Length of dendrites varies from a few ten micrometers to a few hundred micrometers. Detachment and fragmentation of secondary dendrite arms in some areas are also noticeable. XRD spectra of as-cast Ni3Al compound confirms that the microstructure is completely Ni3Al. 3.2. Microstructure of the joints Fig. 4 shows OM images of the microstructure of the interface,
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Materials Today Communications 21 (2019) 100619
M. Soltani Samani, et al.
Fig. 5. Effects of holding time for a) 30, b) 90, and c) 120 min on the microstructure and d) the width of the joint.
the interface has to do with the constant diffusion of alloying elements towards and outwards the interface. This will further be elucidated by SEM/EDS analyses later on this paper. Another worth noting feature is the appearance of sharp corners at the crossover of grain boundaries with the TLP/base metal interface. Examples are highlighted in Fig. 4b. This must have been obvious by the fact that grain boundaries provide easier diffusion path, resulting in the further advancement of the melt towards the base metal. The sharpness of these features decreases in the course of TLP holding time, such a way that after holding for 90 and 120 min, many of these features can be considered round, instead of being sharp (see highlighted area in Fig. 4d). Fig. 5 shows SEM images of the microstructure of the joint, welded with different holding times (30, 90, and 120 min). In some areas, solidification cracking is observed alongside the joint centerline. Increasing the holding time decreases the continuity of the crack at the interface. Solidification cracking is more influenced by the presence of aluminum in the melting pool at the centerline. Aluminum has an inherent tendency to shrink, inducing tensile stresses in the centerline. In case there is not enough melt to fill the shrinking area, cracking at the centerline is expected. As can be seen, the structure at the interface is a typical ASZ structure. Depending
Table 1 EDS analyses of chemical compositions of different areas at the TLP joint at 30 min holding time (areas of analyses are shown in Fig. 5). Area (Wt%)
Ni
Al
Cr
Fe
Si
A B C D
5.2 85.5 81.2 79.9
– 6.7 6.8 10.8
93.65 3.9 4.7 2.2
0.04 1.9 3.0 1.6
0.11 2.0 4.2 5.5
joined at 1050 °C, with holding times ranging from 30 to 120 min. Observed microstructure is a typical athermal solidification zone (ASZ) structure. The dominant feature at the interface is the presence of a network of blocky phases at the middle of the joint. With increasing the holding time, the area fraction of this blocky phase is significantly reduced. It appears that with increasing holding time up to 90 min, the middle blocky phase network becomes more continuous, such a way that after 90 min, no individual block is discernible. Further holding up to 120 min is associated with the formation of a disconnected chain of blocky phases. The variation in the morphology of observed phases at
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Fig. 6. EDS line scan across the interface for TLP joint, with holding time a) 30, b) 60, c) 90, and d) 120 min.
been diffused towards the base metal and from the base metal towards the interlayer. At the onset of solidification, nucleation starts from the base metal and the solidification front moves towards the centerline. While solidification front moves inward the centerline, the solubility of the elements in the melt has decreased, and this in turn is associated with rejection of extra amount of elements to the joint centerline. This obviously decreases the temperature, making the condition ready for a non-equilibrium solidification. Therefore, one can conclude that the holding time in this case is not sufficient to form an isothermal solidification zone. What has been seen at the interface is some possibly brittle and hard phases. Results also show that the microstructure in the vicinity of the solidified interface (the so called DAZ area) is not really affected by the TLP, i.e. there is no indication of precipitation in this area. When it comes to controlling properties in TLP, it is extremely important to have an understanding of how alloying elements diffuse during the process. Alloying elements are obviously distributed between the solidifying front and the melt according to their distribution coefficient. The distribution coefficient (K) of an element is defined as K = CSi/CiL, where CS is the concentration of the solute in the solid and CL is that in the melt. In case the element has a value K < 1, it tends to remain in the melt, making the melt rich in that specific alloy. Thus, in case K < 1, the concentration of alloying elements in front of the solidifying interface increases due to the rejection of alloying element from the solid [12]. Fig. 5d shows the effects of holding time on the width of the joint, showing that there is almost a linear relationship between the width of the joint and the holding time, with the width of the joint changing from approximately 57 μm after 30 min of holding to roughly 75 μm after 120 min. The width of the joint is controlled by the
Fig. 7. XRD spectra of the centerline area of TLP joints with different holding times.
on the chemistry and processing parameters, three typical regions at a TLP joint can be anticipated. The first region is known as athermal solidification zone (ASZ), which is often formed at the centerline. The second region is the so called isothermal solidification zone (ISZ) in which solidification is carried out at a constant temperature. The next region is named diffusion affected zone (DAZ), which as the name suggest is caused by the inter-diffusion of alloying elements from the interlayer towards the base metal. Results show that in this case holding at 1050 °C up to 120 min is not sufficient to achieve an ISZ structure at the centerline of the interface. It appears that the AWS BNi-2 layer has been melted during TLP, and in the meantime alloying elements have
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centerline is associated with an increase in the Si content. Areas C and D have a rather high concentrations of Si, with the former being 4.2 and the latter being 5.5 wt%. It appears that Si is rejected from the centerline, resulting in the formation of Si-rich areas in the vicinity of centerline. It can be postulated that Si cannot dissolve in Ni-Al phase at the centerline and will therefore be rejected outwards the interface. Fig. 6 shows the results of EDS line scan across the interface at different holding times. Increasing the holding time is associated with the diffusion of alloying elements of the interlayer (Cr, Si, and Fe) towards the base metal. This must have been caused by the gradient of concentration between the interlayer and the base metal. Aluminum, on the other hand, has an opposite behavior, such a way that the centerline becomes enriched with Al with increasing the holding time. This obviously is related to the fact that the interlayer does not contain Al, and this provides a gradient of concentration towards the joint centerline. One can also see that the variation of Ni across the interface with holding time is almost negligible, on the grounds that the interlayer contains a large amount of Ni. Fig. 7 depicts XRD spectra of the centerline of TLP joints with different holding times. Results show that different intermetallic compounds, including Cr2B, Ni3Si, and Ni3B are formed at the interface.
Fig. 8. Hardness profilometry across the interface for TLP joint, with holding time a) 30, b) 60, c) 90, and d) 120 min (standard deviation of measurement is 5%).
3.3. Mechanical properties of the joints Fig. 8 shows the hardness variation across the TLP joint, in samples with different holding time (30, 60, 90 and 120 min). Results show that the longer the holding time, the higher hardness can be achieved at the centerline. The increase in hardness is attributable to the enrichment of the centerline from boron and the formation of hard and brittle borides at the centerline. The maximum hardness increase is attained when holding time is increased from 60 to 90 min. Further holding up 120 min has marginal influence on the hardness. Moving away from the centerline in all cases is associated with a rather sharp drop of hardness up to 100 μm from the centerline. It appears that the distance up to 100 μm is the area in which diffusion is controlling the properties of the joint. Results of shear test are shown in Fig. 9. Results are in accordance with the variation of hardness, i.e. the higher the holding time, the higher is the shear strength. Maximum shear strength (219 MPa) is attained after 120 min of holding. Increasing the shear strength by increasing the bonding time at a constant temperature using the AWS BNi-2 interlayer has already been reported by other researchers [11,14,17]. The same arguments for the increase in hardness applies in here as well. Fracture surfaces of TLP joints with holding times 30, 60, 90, and 120 min are shown in Fig. 10. Brittle cleavage fracture (CF) and river pattern (RP) are obviously dominant features in all fracture surfaces. The cleavage area is larger in samples with holding times 30 and 60 min, compared to samples with holding times 90 and 120 min, inferring that these two samples are comparatively more brittle. The observed brittleness is not a surprise, given that TLP joints in all cases have solidification cracking and brittle intermetallic phases at the centerline. As shown earlier in this paper, the centerline zones in TLP joints are mainly composed of chromium borides (Cr2B) and nickel borides (Ni3B), while both are known as inherently very brittle.
Fig. 9. Shear strength of TLP joints, with holding time a) 30, b) 60, c) 90, and d) 120 min.
diffusion of alloying elements. This is best shown by X = K√Dt equation, where X is the diffusion distance of the diffusing element, K is a constant number, D is the diffusion coefficient, and t is the holding time. Based on this equation, the higher the holding time, the larger is the diffusion distance, inferring that elements like Si and B will diffuse further away from the centerline towards the base metal. As mentioned earlier, both Si and B can drastically decrease the melting point. With that said, one can conclude that a decrease in the melting point results in further dissolution/melting of the base metal, and this in turn implies a thicker joint at the end of TLP. Results of EDS analyses of phases (areas), shown in Fig. 5 are given in Table 1. The black blocky phase, occasionally observed at the center of the joint (Area A), is rich in Cr and B. This phase is possibly a chromium boride intermetallic compound. The formation of this chromium boride intermetallic at TLP joints with BNi-2 interlayer has already been reported in the literature [10,11,13–15]. Formation of this phase at the interface indicates that boron is rejected to the center of the joint. This is not surprising, given that boron has an elemental distribution coefficient less than one (k < 1), inferring that boron tends to stay in the melting zone. Other than that boron has a rather low solubility in nickel (approximately 0.3 at.%) [16], supporting the argument that boron tends to segregate to the melting pool. The continuous phase at the center of the joint (area B) is rich in Ni and Al, implicating that the dominant phase at the interface is possibly Ni3Al compound. Further movements outwards the
4. Conclusions Microstructure and mechanical properties of Ni3Al TLP joints, using the AWS BNi-2 interlayer, held at 1050 °C, for 30, 60, 90, and 120 min, were investigated in this study. The following conclusions can be drawn:
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Materials Today Communications 21 (2019) 100619
M. Soltani Samani, et al.
Fig. 10. Fracture surfaces of TLP joints, with holding time a) 30, b) 60, c) 90, and d) 120 min.
• There is no indication of the formation of isothermal solidification •
• • • •
strength. The shear strength increases from 180 to 219 MPa.
• Fractography of fracture surfaces show that in all samples, fracture
zone (ISZ), in neither of samples. The microstructure at the interface in all samples is a typical athermal solidification zone (ASZ). This infers that achieving an ISZ microstructure necessitates a longer holding time. Increasing the holding time is associated with an increase in the width of the joint. There is almost a linear relationship between the width of the joint and the holding time, with the width of the joint changing from approximately 57 μm after 30 min of holding to roughly 75 μm after 120 min. Also, increasing the holding time is accompanied by inter-diffusion of alloying elements towards and outwards the centerline. In all samples, a rather continuous crack is observed at in the middle of the joint. These cracks are typical solidification cracking. As well, the centerline is rich in boron and some boride intermetallic compounds. The maximum hardness in all samples is attained at the centerline, which is in accordance with the arguments of boron segregation to the centerline, followed by the formation of brittle boride intermetallics at the centerline. Increasing the holding time results in an increase in maximum hardness. The maximum hardness 772 HV is attained in sample with holding time 120 min. Increasing the holding time also results in an increase in the shear
mode is brittle, which obviously is related to the fact that centerline contains brittle intermetallic compounds and solidification cracking.
References [1] A. Bahrami, A. Mohammadnejad, M. Sajadi, M. Yazdan Mehr, Spark plasma sintering of Ni3Al-xB-1wt% CNT (0.0 < x < 1.5 at%) nanocomposite, J. Alloys Compd. 788 (2019) 461–467. [2] A. Mohammadnejad, A. Bahrami, M. Sajadi, P. Karimi, H.R. Fozveh, M. Yazdan Mehr, Microstructure and mechanical properties of spark plasma sintered nanocrystalline Ni3Al-xB (0.0 < x < 1.5 at%) alloy, Mater. Today Commun. 17 (2018) 161–168. [3] S. Zhang, T. Lin, P. He, N. Zhao, M. Huang, K.W. Paik, A study on the bonding conditions and nonconductive filler contents on cationic epoxy-based Sn–58Bi solder ACFs joints for reliable flex-on-board applications, IEEE Trans. Compon. Packag. Manuf. Technol. 7 (12) (2017) 2087–2094. [4] M. Yang, Y.H. Ko, J. Bang, T.S. Kim, C.W. Lee, S. Zhang, M. Li, Growth inhibition of interfacial intermetallic compounds by pre-coating oriented Cu6Sn5 grains on Cu substrates, J. Alloys Compd. 701 (2017) 533–541. [5] S. Zhang, T. Lin, P. He, K.W. Paik, Effects of acrylic adhesives property and optimized bonding parameters on Sn58Bi solder joint morphology for flex-on-board assembly, Microelectron. Reliab. 78 (2017) 181–189. [6] D.Y. Zhang, H.Y. Gao, K. Lu, H.L. Luo, S.P. Li, Ni3Al-base alloy of laser welding behavior, Adv. Mater. Res. 418 (2012) 1531–1537. [7] M.L. Santella, Weld solidification cracking in cast Ni3Al alloys, Scr. Metall. Mater. 28 (1993) 1305–1310. [8] D.H. Jung, A. Sharma, M. Mayer, J.P. Jung, A review on recent advances in
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[9] [10] [11] [12]
transient liquid phase (TLP) bonding for thermoelectric power module, Rev. Adv. Mater. Sci. 53 (2018) 147–160. K. Bai, L.N. Fern, L. Teck, T.L. Tan, D. Pan, Understanding non-parabolic solidification kinetics in Ni-based alloys during TLP bonding via thermo-kinetic modelling, J. Alloys Compd. 699 (2017) 1084–1094. O.A. Ojo, A. Ghoneim, J. Hunedy, Transient liquid phase bonding of single crystal Ni3Al-based intermetallic alloy, Int. Brazing and Soldering Conference American Society for Metals (2012). T.S. Lin, H.X. Li, H. Peng, Y.A.N.G. Xu, Y.D. Huang, L.I. Liang, H.A.N. Leng, Effect of bonding parameters on microstructures and properties during TLP bonding of Nibased super alloy, Trans. Nonferrous Met. Soc. China 22 (2012) 2112–2117. M.H. Avazkonandeh-Gharavol, M. Haddad-Sabzevar, H. Fredriksson, Effect of partition coefficient on microsegregation during solidification of aluminium alloys, Int. J. Miner. Metall. Mater. 21 (2014) 980–989.
[13] A. Ghoneim, O.A. Ojo, On the influence of boron-addition on TLP bonding time in a Ni3Al-based intermetallic, Intermetallics 18 (2010) 582–586. [14] R. Abdolvand, M. Atapour, M. Shamanian, A. Allafchian, The effect of bonding time on the microstructure and mechanical properties of transient liquid phase bonding between SAF 2507 and AISI 304, J. Manuf. Process. 25 (2017) 172–180. [15] A. Shamsabadi, A. Yarmou, R. Bakhtiari, TLP bonding of IN738/MBF20/IN718 system, J. Alloys Compd. 685 (2016) 896–904. [16] M. Pouranvari, A. Ekrami, A.H. Kokabi, Solidification and solid state phenomena during TLP bonding of IN718 superalloy using Ni–Si–B ternary filler alloy, J. Alloys Compd. 563 (2013) 143–149. [17] V. Jalilvand, H. Omidvar, H.R. Shakeri, M.R. Rahimipour, A study on the effect of process parameters on the properties of joint in TLP-bonded Inconel 738LC superalloy, Metall. Mater. Trans. B 44 (2013) 1222–1231.
8
Contents lists available at ScienceDirect
Materials Today Communications journal homepage: www.elsevier.com/locate/mtcomm
Microstructure and mechanical properties of transient liquid phase (TLP)bonded Ni3Al intermetallic compounds
T
M. Soltani Samani, A. Bahrami , F. Karimzadeh ⁎
Department of Materials Engineering, Isfahan University of Technology, Isfahan 84156-83111, Iran
ARTICLE INFO
ABSTRACT
Keywords: Transient liquid phase bonding (TLP) Ni3Al Joining Intermetallic AWS BNi-2 Microstructure
This paper investigates microstructure and mechanical properties of transient-liquid-phase (TLP) bonding for the joining of Ni3Al intermetallic compound. The jointing was carried by using AWS BNi-2 interlayer in a vacuum furnace at 1050 °C. Effects of holding times for 30, 60, 90, and 120 min on the microstructure and mechanical properties of the joint have been investigated. The microstructure of samples was evaluated using optical and scanning electron microscopes (SEM). Phase identification was carried using energy dispersive spectroscopy (EDS) and X-ray diffractometer (XRD). Mechanical properties of joints were studied by hardness profilometry across the interface as well as shear test. Results show that no isothermally solidified zone has been formed in holding time up to 120 min. Also, solidification cracking was observed in the centerline of joints. EDS analysis showed that the centerline is rich in boron and some boride intermetallic compounds. Increasing holding time is associated with a decrease in the amount of brittle phases in the athermally solidified zone. Microhardness profilometry across the interface show that the maximum hardness is attained in the middle of the joint. This is attributable to the fact that the center of the joint is rich in hard nickel and chromium boride phases. Results also show that both hardness and shear strength of the joint increase with increasing holding time. Fractography results indicate that fracture takes place in a brittle mode.
1. Introduction Intermetallic compounds are known as materials with unique characteristics, including high temperature strength, high temperature oxidation/corrosion resistance, and excellent structural stability at high temperatures. This makes these alloys excellent candidates for many demanding applications in advanced industries. Amongst these intermetallic compounds, Ni3Al intermetallic compound is one of the most widely used ones, owing to the fact that it has a comparatively better formability. Ni3Al intermetallic compound is currently used as a commercial alloy in compressors and turbine blades in different industries [1,2]. A major obstacle in further using of this intermetallic alloy is the welding/joining of this compound. This difficulty originates from the inherent brittleness of intermetallic compounds. Obviously, amongst different fusion and solid state-based welding techniques [3–5], those which induce lower residual stress and heat input in the weldment are more preferred. Fusion-based welding processes, including gas tungsten arc welding (GTAW), electron beam welding (EBW) and laser beam welding are intrinsically prone to solidification-induced defects (cavities, solidification cracking, incomplete welds, voids, …). Given that intermetallic alloys are mostly used in very sensitive and demanding ⁎
applications, hardly is there any tolerance for mentioned defects, inferring that fusion-based techniques must be used with extreme precautions. In an attempt to minimize solidification cracking in laserwelded Ni3Al alloy, Zhang and et al. [6] postulated that controlling the heat input is a key controlling issue, when it comes to the sensitivity of cracks in the heat affect zone (HAZ) area. Obviously, the coarse-grain structure is more sensitive to solidification cracking during welding. Santla and et al. [7] did the same study during the GTAW process. They postulated that controlling the chemistry of the alloy is very influential in minimizing solidification cracking. Contrary to fusion welding techniques, solid-state-based methods, like friction welding process, have much less heat inputs. On the other hand, the downside of friction welding is that the workpiece undergoes excessive plastic deformation, making this method less applicable for the obvious reason of brittleness of intermetallic compounds. With that said, transient liquid-phase process (TLP) appears to be very promising. TLP, known to be a combination of diffusion bonding and brazing, is based on the formation of a transitional melting phase at the interface. Diffusion of alloying elements across the interface during TLP, while temperature is kept constant, ultimately results in the formation of the joint [8]. Microstructure and mechanical properties of TLP joints largely depend on the type and
Corresponding author. E-mail address: [email protected] (A. Bahrami).
https://doi.org/10.1016/j.mtcomm.2019.100619 Received 28 May 2019; Received in revised form 25 August 2019; Accepted 26 August 2019 Available online 27 August 2019 2352-4928/ © 2019 Elsevier Ltd. All rights reserved.
Materials Today Communications 21 (2019) 100619
M. Soltani Samani, et al.
Fig. 1. Schematics of the holder for the shear test.
Fig. 2. Dendritic structure of as-cast Ni3Al compound.
the chemical composition of the interlayer. Normally, melting point depressant elements (e.g. boron) are present in the interlayer to reduce the melting point of the interlayer. The lower the melting point, the easier is the formation of isothermal solidification zone at the joint [9]. Ojo and et al. [10] joined Ni3Al-based single crystal intermetallic alloy using TLP. They showed that insufficient holding time results in the appearance of undesirable brittle eutectic compounds at the joint centerline. They postulated that a mixture of brazing composite powder and interlayer can be used to reduce the time to complete isothermal solidification. However, their findings are limited to single crystals. To our knowledge, hardly is there any comprehensive study on the microstructure and mechanical properties of TLP-joined polycrystalline Ni3Al alloys. This paper investigates the TLP joining of polycrystalline Ni3Al intermetallic compound, using AWS BNi-2 as interlayer. AWS BNi-2 contains B and Si, with both are known to be melting point depressant elements. Besides, AWS BNi-2 has reportedly a very good corrosion resistance and creep strength [11]. In this study, the effects of holding time at a constant temperature/pressure on the mechanical properties and microstructure of bonding are also investigated. Hardly is there any comprehensive report on the microstructure and mechanical properties of TLP-bonded Ni3Al intermetallic compounds in the literature. This paper is an attempt to investigate how properties of TLPbonded Ni3Al intermetallic compounds can be controlled by changing the microstructure. 2. Experimental procedure
Fig. 3. XRD spectrum of as-cast Ni3Al compound, in which peaks coincide with Ni3Al peaks.
The Ni3Al intermetallic compound was first synthesized, in a
2
Materials Today Communications 21 (2019) 100619
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Fig. 4. Effects of holding time for a) 30, b) 60, c) 90, and d) 120 min on the microstructure of the joint.
vacuum arc remelting (VAR) furnace, using pure nickel (99%) and aluminum rod (99.9%). The VAR was carried out three times on each specimen to make sure that alloying elements were homogeneously mixed. Experimental samples for joining were cut in 3 mm thick 10 × 15 mm pieces using electro discharge machining (EDM). TLP was performed in YMVR 1700 vacuum furnace under vacuum of 10−5 bar. The 50 μm thick interlayer AWS BNi-2, used in this study, had the following composition: Ni-5.9Cr-7.0Si-2.8Fe-3B (wt.%). The TLP was performed at 1050 °C four different holding times: 30, 60, 90, and 120 min. Samples were first ground before joining with polishing paper 1200. Both the interlayer and samples were sonicated in acetone for 25 min to ensure that surfaces were not contaminated. A fixture with dimensions of 10 × 20 × 70 mm was used to hold samples together during TLP. Microstructural studies were carried out using optical and electron microscopies. Samples were first ground with grinding papers, followed by polishing with SiC polishing paste. Marble solution (0.5 HCl+H2O+0.5 CuSO4) was used to etch samples. Energy dispersive XRay spectroscopy (EDS) was used to identify phases in the microstructure and to do the line scan across the interface. Vickers hardness profilometry was conducted across the interface, using 100 g load. To
determine the shear strength of the joint, samples of 10 × 10 mm were cut and placed in a fixture, as schematically shown in Fig. 1, then they were then loaded until they broke. The loading speed was chosen to be 1 mm/min. Finally, scanning electron microscope (SEM) was used to do the fractography on the fracture surface (Fig. 2). 3. Results and discussion 3.1. Microstructure of base metal Fig. 3 shows the microstructure of as-cast Ni3Al compound, in which a dendritic structure is clearly visible. Length of dendrites varies from a few ten micrometers to a few hundred micrometers. Detachment and fragmentation of secondary dendrite arms in some areas are also noticeable. XRD spectra of as-cast Ni3Al compound confirms that the microstructure is completely Ni3Al. 3.2. Microstructure of the joints Fig. 4 shows OM images of the microstructure of the interface,
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Fig. 5. Effects of holding time for a) 30, b) 90, and c) 120 min on the microstructure and d) the width of the joint.
the interface has to do with the constant diffusion of alloying elements towards and outwards the interface. This will further be elucidated by SEM/EDS analyses later on this paper. Another worth noting feature is the appearance of sharp corners at the crossover of grain boundaries with the TLP/base metal interface. Examples are highlighted in Fig. 4b. This must have been obvious by the fact that grain boundaries provide easier diffusion path, resulting in the further advancement of the melt towards the base metal. The sharpness of these features decreases in the course of TLP holding time, such a way that after holding for 90 and 120 min, many of these features can be considered round, instead of being sharp (see highlighted area in Fig. 4d). Fig. 5 shows SEM images of the microstructure of the joint, welded with different holding times (30, 90, and 120 min). In some areas, solidification cracking is observed alongside the joint centerline. Increasing the holding time decreases the continuity of the crack at the interface. Solidification cracking is more influenced by the presence of aluminum in the melting pool at the centerline. Aluminum has an inherent tendency to shrink, inducing tensile stresses in the centerline. In case there is not enough melt to fill the shrinking area, cracking at the centerline is expected. As can be seen, the structure at the interface is a typical ASZ structure. Depending
Table 1 EDS analyses of chemical compositions of different areas at the TLP joint at 30 min holding time (areas of analyses are shown in Fig. 5). Area (Wt%)
Ni
Al
Cr
Fe
Si
A B C D
5.2 85.5 81.2 79.9
– 6.7 6.8 10.8
93.65 3.9 4.7 2.2
0.04 1.9 3.0 1.6
0.11 2.0 4.2 5.5
joined at 1050 °C, with holding times ranging from 30 to 120 min. Observed microstructure is a typical athermal solidification zone (ASZ) structure. The dominant feature at the interface is the presence of a network of blocky phases at the middle of the joint. With increasing the holding time, the area fraction of this blocky phase is significantly reduced. It appears that with increasing holding time up to 90 min, the middle blocky phase network becomes more continuous, such a way that after 90 min, no individual block is discernible. Further holding up to 120 min is associated with the formation of a disconnected chain of blocky phases. The variation in the morphology of observed phases at
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Fig. 6. EDS line scan across the interface for TLP joint, with holding time a) 30, b) 60, c) 90, and d) 120 min.
been diffused towards the base metal and from the base metal towards the interlayer. At the onset of solidification, nucleation starts from the base metal and the solidification front moves towards the centerline. While solidification front moves inward the centerline, the solubility of the elements in the melt has decreased, and this in turn is associated with rejection of extra amount of elements to the joint centerline. This obviously decreases the temperature, making the condition ready for a non-equilibrium solidification. Therefore, one can conclude that the holding time in this case is not sufficient to form an isothermal solidification zone. What has been seen at the interface is some possibly brittle and hard phases. Results also show that the microstructure in the vicinity of the solidified interface (the so called DAZ area) is not really affected by the TLP, i.e. there is no indication of precipitation in this area. When it comes to controlling properties in TLP, it is extremely important to have an understanding of how alloying elements diffuse during the process. Alloying elements are obviously distributed between the solidifying front and the melt according to their distribution coefficient. The distribution coefficient (K) of an element is defined as K = CSi/CiL, where CS is the concentration of the solute in the solid and CL is that in the melt. In case the element has a value K < 1, it tends to remain in the melt, making the melt rich in that specific alloy. Thus, in case K < 1, the concentration of alloying elements in front of the solidifying interface increases due to the rejection of alloying element from the solid [12]. Fig. 5d shows the effects of holding time on the width of the joint, showing that there is almost a linear relationship between the width of the joint and the holding time, with the width of the joint changing from approximately 57 μm after 30 min of holding to roughly 75 μm after 120 min. The width of the joint is controlled by the
Fig. 7. XRD spectra of the centerline area of TLP joints with different holding times.
on the chemistry and processing parameters, three typical regions at a TLP joint can be anticipated. The first region is known as athermal solidification zone (ASZ), which is often formed at the centerline. The second region is the so called isothermal solidification zone (ISZ) in which solidification is carried out at a constant temperature. The next region is named diffusion affected zone (DAZ), which as the name suggest is caused by the inter-diffusion of alloying elements from the interlayer towards the base metal. Results show that in this case holding at 1050 °C up to 120 min is not sufficient to achieve an ISZ structure at the centerline of the interface. It appears that the AWS BNi-2 layer has been melted during TLP, and in the meantime alloying elements have
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centerline is associated with an increase in the Si content. Areas C and D have a rather high concentrations of Si, with the former being 4.2 and the latter being 5.5 wt%. It appears that Si is rejected from the centerline, resulting in the formation of Si-rich areas in the vicinity of centerline. It can be postulated that Si cannot dissolve in Ni-Al phase at the centerline and will therefore be rejected outwards the interface. Fig. 6 shows the results of EDS line scan across the interface at different holding times. Increasing the holding time is associated with the diffusion of alloying elements of the interlayer (Cr, Si, and Fe) towards the base metal. This must have been caused by the gradient of concentration between the interlayer and the base metal. Aluminum, on the other hand, has an opposite behavior, such a way that the centerline becomes enriched with Al with increasing the holding time. This obviously is related to the fact that the interlayer does not contain Al, and this provides a gradient of concentration towards the joint centerline. One can also see that the variation of Ni across the interface with holding time is almost negligible, on the grounds that the interlayer contains a large amount of Ni. Fig. 7 depicts XRD spectra of the centerline of TLP joints with different holding times. Results show that different intermetallic compounds, including Cr2B, Ni3Si, and Ni3B are formed at the interface.
Fig. 8. Hardness profilometry across the interface for TLP joint, with holding time a) 30, b) 60, c) 90, and d) 120 min (standard deviation of measurement is 5%).
3.3. Mechanical properties of the joints Fig. 8 shows the hardness variation across the TLP joint, in samples with different holding time (30, 60, 90 and 120 min). Results show that the longer the holding time, the higher hardness can be achieved at the centerline. The increase in hardness is attributable to the enrichment of the centerline from boron and the formation of hard and brittle borides at the centerline. The maximum hardness increase is attained when holding time is increased from 60 to 90 min. Further holding up 120 min has marginal influence on the hardness. Moving away from the centerline in all cases is associated with a rather sharp drop of hardness up to 100 μm from the centerline. It appears that the distance up to 100 μm is the area in which diffusion is controlling the properties of the joint. Results of shear test are shown in Fig. 9. Results are in accordance with the variation of hardness, i.e. the higher the holding time, the higher is the shear strength. Maximum shear strength (219 MPa) is attained after 120 min of holding. Increasing the shear strength by increasing the bonding time at a constant temperature using the AWS BNi-2 interlayer has already been reported by other researchers [11,14,17]. The same arguments for the increase in hardness applies in here as well. Fracture surfaces of TLP joints with holding times 30, 60, 90, and 120 min are shown in Fig. 10. Brittle cleavage fracture (CF) and river pattern (RP) are obviously dominant features in all fracture surfaces. The cleavage area is larger in samples with holding times 30 and 60 min, compared to samples with holding times 90 and 120 min, inferring that these two samples are comparatively more brittle. The observed brittleness is not a surprise, given that TLP joints in all cases have solidification cracking and brittle intermetallic phases at the centerline. As shown earlier in this paper, the centerline zones in TLP joints are mainly composed of chromium borides (Cr2B) and nickel borides (Ni3B), while both are known as inherently very brittle.
Fig. 9. Shear strength of TLP joints, with holding time a) 30, b) 60, c) 90, and d) 120 min.
diffusion of alloying elements. This is best shown by X = K√Dt equation, where X is the diffusion distance of the diffusing element, K is a constant number, D is the diffusion coefficient, and t is the holding time. Based on this equation, the higher the holding time, the larger is the diffusion distance, inferring that elements like Si and B will diffuse further away from the centerline towards the base metal. As mentioned earlier, both Si and B can drastically decrease the melting point. With that said, one can conclude that a decrease in the melting point results in further dissolution/melting of the base metal, and this in turn implies a thicker joint at the end of TLP. Results of EDS analyses of phases (areas), shown in Fig. 5 are given in Table 1. The black blocky phase, occasionally observed at the center of the joint (Area A), is rich in Cr and B. This phase is possibly a chromium boride intermetallic compound. The formation of this chromium boride intermetallic at TLP joints with BNi-2 interlayer has already been reported in the literature [10,11,13–15]. Formation of this phase at the interface indicates that boron is rejected to the center of the joint. This is not surprising, given that boron has an elemental distribution coefficient less than one (k < 1), inferring that boron tends to stay in the melting zone. Other than that boron has a rather low solubility in nickel (approximately 0.3 at.%) [16], supporting the argument that boron tends to segregate to the melting pool. The continuous phase at the center of the joint (area B) is rich in Ni and Al, implicating that the dominant phase at the interface is possibly Ni3Al compound. Further movements outwards the
4. Conclusions Microstructure and mechanical properties of Ni3Al TLP joints, using the AWS BNi-2 interlayer, held at 1050 °C, for 30, 60, 90, and 120 min, were investigated in this study. The following conclusions can be drawn:
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Fig. 10. Fracture surfaces of TLP joints, with holding time a) 30, b) 60, c) 90, and d) 120 min.
• There is no indication of the formation of isothermal solidification •
• • • •
strength. The shear strength increases from 180 to 219 MPa.
• Fractography of fracture surfaces show that in all samples, fracture
zone (ISZ), in neither of samples. The microstructure at the interface in all samples is a typical athermal solidification zone (ASZ). This infers that achieving an ISZ microstructure necessitates a longer holding time. Increasing the holding time is associated with an increase in the width of the joint. There is almost a linear relationship between the width of the joint and the holding time, with the width of the joint changing from approximately 57 μm after 30 min of holding to roughly 75 μm after 120 min. Also, increasing the holding time is accompanied by inter-diffusion of alloying elements towards and outwards the centerline. In all samples, a rather continuous crack is observed at in the middle of the joint. These cracks are typical solidification cracking. As well, the centerline is rich in boron and some boride intermetallic compounds. The maximum hardness in all samples is attained at the centerline, which is in accordance with the arguments of boron segregation to the centerline, followed by the formation of brittle boride intermetallics at the centerline. Increasing the holding time results in an increase in maximum hardness. The maximum hardness 772 HV is attained in sample with holding time 120 min. Increasing the holding time also results in an increase in the shear
mode is brittle, which obviously is related to the fact that centerline contains brittle intermetallic compounds and solidification cracking.
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