**6. Mechanism of joint formation**

The mechanism of bond formation using composite Ni-Al2O3 coating is discussed in this section. The experimental results and the scientific literature show that the joint formation during transient liquid phase diffusion bonding is completed in five distinct stages: interfacial contact and solid-state diffusion, eutectic melting and base metal dissolution, and isothermal solidification. These stages will be discussed thoroughly with reference to the change in composition across the joint region. Mathematical models for predicting the parameter settings when nano-composite coatings are used in joining will also be presented. The sub-section of this topic are as follows:

## **6.1. Interfacial contact and Solid-state diffusion**

330 Advanced Aspects of Spectroscopy

in strengthening the nano-composites.

μm thick Ni–(50nm) Al2O3 for 10 minutes at 600oC.

**6. Mechanism of joint formation** 

The sub-section of this topic are as follows:

MPa with 500 nm to142 MPa with 50 nm. This increase in joint shear strength was attributed to better distribution of nano-sized particles within the interlayer when smaller particle sizes are used. In both cases higher shear strengths were obtained than when pure Ni coating is used (117 MPa) [23]. The results indicate that joint strength of up to 90% that of the base metal (BM) is achievable when using a 50 nm diameter nano-sized particle-reinforced interlayer. Tjong [42] showed that the nano-particle size has a strong effect on the yield strength. The author suggested that a particle size of 100 nm is a critical value for improving the yield strength of nano-composites. Below this critical value the yield strength increases significantly with decreasing particle size. Similar results were obtained by Gupta and coworkers [43, 44]. Zhang and Chen [45] showed that the Orowan stress plays a major role

Figure 21 shows the fractured surface for a bond made using a 5μm Ni- 50 nm Al2O3 particle. The fractograph showed shear plastic deformation, indicative of ductile fracture with a crack propagating primarily through the bond-line and a section of the base metal adjacent to the bond-line. A similar result was obtained when a dispersed particle size of 50 nm were used in the coating. XRD analyses of the fractured surfaces indicated the presence

**Figure 21.** (a) SEM micrograph and (c) XRD analysis of the fractured surface for a bond made with 5

The mechanism of bond formation using composite Ni-Al2O3 coating is discussed in this section. The experimental results and the scientific literature show that the joint formation during transient liquid phase diffusion bonding is completed in five distinct stages: interfacial contact and solid-state diffusion, eutectic melting and base metal dissolution, and isothermal solidification. These stages will be discussed thoroughly with reference to the change in composition across the joint region. Mathematical models for predicting the parameter settings when nano-composite coatings are used in joining will also be presented.

of peaks for Al2O3, NiAl2O4 and Al11Ni9 compound at the fractured surface.

The first stage of transient liquid phase bonding involved heating the sample to the bonding temperature. During this stage of bonding, two mechanisms are thought to occur simultaneously, firstly an increase in interfacial contact between the coating surface and the Al 6061 surface and secondly, solid-state diffusion along the coating/MMC contact interface. The initial contact area between the one side of the coating and the metal surfaces is only a small fraction of the theoretical area available, due to the presences of micro-asperities of the surface of both the metal sample and the coating. However, under the effects of heating and an external pressure an intimate contact can be establish at the bonding surfaces, as the micro-asperities suffer plastic deformation. As the temperature increases greater plastic flow is achieved at the interface and the percentage contact area increases. This increase in the contact area results in an increased diffusivity of the solute (Ni) into the base metal. During the heating stage Ni diffuses deep into the Al-MMCs resulting in the formation of complex intermetallic compounds.

Figure 22(a) shows an SEM micrograph of a joint bonded using a 15μm thick Ni-Al2O3 coating as the interlayer. The joint was made at a bonding temperature of 600oC for 1 minute and shows three distinct reaction layers at the interlayer/MMC interface. The nano- Al2O3 particles that were co-electrodeposited with Ni can are clearly shown in Figure 22.

The composition of the reaction layers was determined quantitatively by energy dispersive spectroscopy (EDS) analysis and is shown in Table 4. The formation of reaction layers (L2 and L3) shown in Figure 22, occurred as a result of the inter-diffusion of Ni and Al. EDS analysis showed that L1 was composed of a Ni-Al layer dispersed with nano-sized Al2O3 particles after a bonding time of 1 minute (see Table 4). Reaction layer L2 on the other hand appears to be a nickel-aluminide with compositions of 50.7 (at %) Ni and 46.70(at %) Al. The L3 layer contains approximately 24.30 (at %) Ni and 77.0 (at %) Al, which is likely to be NiAl3 intermetallic. This compound is believed to form due to the low solubility of Ni in Al. This has been reported to be approximately 2.9 at% [46]. The saturation of the aluminum interface through the inter-diffusion of Ni and Al leads to the precipitation of the nickel aluminide intermetallic NiAl3.

The phase diagram of the Ni + Al system indicates the thermodynamic stability of the γ'- Ni3Al phase when formed in the nickel concentration range of about (74 to 76) at.% [46]. Additionally, the phase diagram of the (Ni + Al) system proposed by Nash et al. [47] showed that for aluminum concentrations exceeding 40 mol%, there exist three coexistence fields: (Al + NiAl3), (NiAl3 + Ni2Al3) and the non-stoichiometric intermetallic β -NiAl, which is formed in the concentration range 43mol% to 59 mol% aluminum. Rog et al. [48] determined the Gibbs free energy of formation for various intermetallic compounds forming in the Ni + Al system. The results showed that within the temperature range of 570oC to 620oC (843K to 893K) the nickel aluminide compounds listed in Table 4 are formed.

Based on the scientific literature, the compound formed in the reaction layer L3 is likely to be the NiAl3. This compound also appears on the right side of Equation 9 and is believed to

form due to low solubility of Ni in Al which has been reported to be approximately 2.9 at% [46]. The saturation of the aluminum interface through the diffusion of Ni can lead to the precipitation of the nickel aluminide intermetallic NiAl3. The composition of L2 indicates that the compound is likely NiAl.


**Table 4.** The standard Gibbs free energy values for the chemical reactions with nickel aluminides at 870K [48]


**Table 5.** Energy dispersive spectroscopic compositional analyses of the reaction layers developed during bonding (wt%)

**Figure 22.** (a) SEM micrograph of joint bonded with a 15 μm Ni-Al2O3 coating for 1 min.

#### **6.2. Liquid formation and base-metal dissolution**

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870K [48]

during bonding (wt%)

that the compound is likely NiAl.

**Layers Al /**

form due to low solubility of Ni in Al which has been reported to be approximately 2.9 at% [46]. The saturation of the aluminum interface through the diffusion of Ni can lead to the precipitation of the nickel aluminide intermetallic NiAl3. The composition of L2 indicates

**Phase** o 1 <sup>Δ</sup>G (kJ.mol ) <sup>f</sup>

Ni +Al = NiAl -133.0 +/- 1.0 Ni+NiAl3= Ni2Al3 -144.1+/- 0.8 2Ni +3Al = Ni2Al3 -311.0+/- 1.7 Ni + 3Al =NiAl3 -166.8+/- 0.9 **Table 4.** The standard Gibbs free energy values for the chemical reactions with nickel aluminides at

wt% **Ni /** wt% **Si /** wt% **Mg /**

**Table 5.** Energy dispersive spectroscopic compositional analyses of the reaction layers developed

**Figure 22.** (a) SEM micrograph of joint bonded with a 15 μm Ni-Al2O3 coating for 1 min.

L1 12.5 87.5 0 0 Ni +Al2O3 L2 46.71 50.67 0.27 0 NiAl L3 72.84 24.31 0 0.55 NiAl3 L4 72.84 24.31 0 0.55 NiAl3

wt% **Compound** 

Immediately following the heating stage, eutectic melting ensues. According to Dmitry et al [50] there are two possible reactions which are capable of producing a liquid at the joint interface within this temperature range (570oC-620oC). Upon reaching a temperature of 565oC, the L2 compound reacts with L4 (88.4%Al and 2.08 % Si) to form a eutectic liquid (E1) along the bond interface as predicted by Equation 9 [46]. The compound formed in L2 is consumed in the eutectic reaction and diffuses into the base metal. When the joint region was heated to 577oC a second eutectic liquid formed within L4 as predicted by Equation 10 [49]. The formation of a eutectic liquid layer at the bond interface leads to faster interdiffusion between Al and the Ni interlayer and this results in a gradual change in the composition of the joint region.

$$\text{Li}\_{1} \xrightarrow{565\,^{\text{O}}\text{C}} \text{(87\%)}\text{Al} + \text{(2\%)}\text{Si} + \text{(11\%)}\text{NiAl}\_{\text{3}}\tag{9}$$

$$\text{Li}\_2 \xrightarrow{577\,^\text{O}\text{C}} \text{(87.5\%)}\text{Al} + \text{(12.5\%)}\text{Si} \tag{10}$$

Dmitry et al. [50] carried out thermodynamic calculations of (Al-Mg-Si-Fe-Ni) quinary systems formed in aluminum alloys. The results showed that in alloys containing Al-Mg-Si-Fe-Ni, numerous ternary and quaternary reactions can occur that has the potential of producing a liquid phase. Some of the phases that contribute to these reactions are shown in Table 6. The XRD analysis of the fractured surfaces shown in the previous also indicated that presence of AlFe3Si. This phase is possibly a variant of the β-phase family listed in Table 6.


**Table 6.** Chemical composition and density of phases formed in Al-Mg-Si-Fe-Ni system [50]

#### **6.3. Base metal dissolution**

Base metal dissolution is also an important part of the second stage of bonding. Research work published [1,2] on the mechanisms of TLP bonding showed that dissolution of the interlayer and base metal occurs simultaneously at the bonding temperature. In this study, holding at a bonding temperature above 577oC resulted in increased diffusion of Ni into the base metal and causes the liquid phase to spread between the bonding surfaces due to the effects of pressure and capillary action. The application of pressure enhances spreading of the liquid phase between the bonding surfaces due to capillary action. This spreading increases the contact area and induces the diffusion of Ni and Al into the liquid phase. This

results in more eutectic liquid formation and an increase in the width of the liquid phase at the joint, due to dissolution of a section of the base metal. Further increase in the bonding time to 5 minutes, resulted in the diffusion of the Ni into the base metal and away from the joint interface leading to the formation of eutectic or peritectic liquid along the grain boundary as indicated in the EDS analysis of the eutectic microstructure shown in Figure 23.

**Figure 23.** (a) SEM micrograph of joint bonded using a 25 μm Ni foil for 5 min and (b) EDS analysis of the joint region.

According to published studies [20,1] on the stages of TLP bonding, it is expected that dissolution of the interlayer and base metal occurs simultaneously at the bonding temperature followed by spreading of the liquid phase between the bonding surfaces. This spreading increased the bonded area and enhanced the diffusion of Ni and Al into the liquid phase. This continuing diffusion resulted in more liquid formation and an increase in the width of liquid phase. The maximum liquid width attained when using the Ni-Al2O3 composite coating as the interlayer was found by Cooke et al [24] to be max 20.6 . *W wo*

## **6.4. Isothermal solidification**

In TLP bonding, prolonged hold time at the bonding temperature for 10 minutes allowed for the diffusion of Al into the eutectic liquid which caused the composition of the liquid phase to become aluminum rich, resulting in a change in the eutectic composition (see Table 6). The change in the joint composition initiates the isothermal solidification stage of TLP bonding as a function of bonding time since the temperature and interlayer thickness is constant. When the bonding time is increased to 10 minutes the interface is eliminated and the grain size within the joint increased. This disrupts the band of segregated particles at the interface and homogenized the joint zone. When the bonding time was increased to 30 minutes a corresponding increase was seen in grain size, resulting in more uniform distribution of micro-Al2O3 particles. In the reported studies on transient liquid phase diffusion bonding of Al-MMCs, it was shown that the width of the segregated zone at the joint center increased with increasing bonding time. The opposite of this relationship was seen when using the Ni-Al2O3 coating. As the bonding time increased, the width of the segregated region decrease. This can be attributed to the heterogeneous nucleation of grains within the joint zone during solidification and this lead to grain refining at the joint. The high resolution SEM micrograph shown in Figure 33 revealed the presence of a nano-Al2O3 particle at the center of a grain. EDX spectra of the particle showed Al and O in high concentrations with traces of Mg. Comparing Gibbs free energy of formation at the bonding temperature for MgO (-1195 kJ/mol) and Al2O3 (-985 kJ/mol) it is found that Al2O3 is unstable in the presence of Mg hence it is like that some of the nano-size Al2O3 will decomposed to form MgAl2O4 compound.

334 Advanced Aspects of Spectroscopy

the joint region.

**6.4. Isothermal solidification** 

results in more eutectic liquid formation and an increase in the width of the liquid phase at the joint, due to dissolution of a section of the base metal. Further increase in the bonding time to 5 minutes, resulted in the diffusion of the Ni into the base metal and away from the joint interface leading to the formation of eutectic or peritectic liquid along the grain boundary as indicated in the EDS analysis of the eutectic microstructure shown in Figure 23.

**Figure 23.** (a) SEM micrograph of joint bonded using a 25 μm Ni foil for 5 min and (b) EDS analysis of

According to published studies [20,1] on the stages of TLP bonding, it is expected that dissolution of the interlayer and base metal occurs simultaneously at the bonding temperature followed by spreading of the liquid phase between the bonding surfaces. This spreading increased the bonded area and enhanced the diffusion of Ni and Al into the liquid phase. This continuing diffusion resulted in more liquid formation and an increase in the width of liquid phase. The maximum liquid width attained when using the Ni-Al2O3 composite coating as the interlayer was found by Cooke et al [24] to be max 20.6 . *W wo*

In TLP bonding, prolonged hold time at the bonding temperature for 10 minutes allowed for the diffusion of Al into the eutectic liquid which caused the composition of the liquid phase to become aluminum rich, resulting in a change in the eutectic composition (see Table 6). The change in the joint composition initiates the isothermal solidification stage of TLP bonding as a function of bonding time since the temperature and interlayer thickness is constant. When the bonding time is increased to 10 minutes the interface is eliminated and the grain size within the joint increased. This disrupts the band of segregated particles at the interface and homogenized the joint zone. When the bonding time was increased to 30 minutes a corresponding increase was seen in grain size, resulting in more uniform distribution of micro-Al2O3 particles. In the reported studies on transient liquid phase diffusion bonding of Al-MMCs, it was shown that the width of the segregated zone at the joint center increased with increasing bonding time. The opposite of this relationship was seen when using the Ni-Al2O3 coating. As the bonding time increased, the width of the The segregation of particles during isothermal solidification was accredited to the pushing of strengthening particles by the solidifying liquid-solid interface. Stefanescu [17, 18] showed that particle pushing can be assumed to be a steady-state condition under which the interface velocity can be assumed to be equal the rate of isothermal solidification This rate can be calculated using a model proposed by Sinclair [51]. The constant, ξ, signifies the solidification rate of the system. Increasing ξ results in faster solid-liquid interface motion, and a shorter duration of the isothermal solidification stage. The rate of isothermal solidification can be calculated using Equation 11.

$$\xi = -2\begin{pmatrix} k \ -1 \end{pmatrix}^{-1} \sqrt{\frac{D}{\pi}} \cdot \frac{\exp\left(\frac{-\xi^2}{4D}\right)}{\operatorname{erfc}\left(\frac{\xi}{2\sqrt{D}}\right)}\tag{11}$$

Where *k* is a partition coefficient given by *<sup>L</sup> L C C* and D is the diffusivity of Ni into Al. The

diffusivity at 570oC is 4.69 x10-13 m2/s, when the bonding temperature is increased to 620oC, the diffusivity also increased to 1.58 x10-12 m2/s. This increase in diffusivity is reflected in a faster solid-liquid interface rate () and a shorter isothermal solidification stage.

The final concentration of Ni *CL* was taken from the Al-Ni-Si phase diagram [30] to be 4.9 wt % for the bonding temperature of 620°C. The diffusivity of Ni in Al at 620°C is D = 1.58 x 10-12 m2/s [25, 31]. Using these values the predicted interface rate constant -0.395μm/s was calculated from Equation 8. This solidification rate is significantly less than the critical interface velocity (16 -400 μm/s) required to engulf dispersed particle during solidification [26]. Li et al. [21] suggested that particle segregation tendency is dependent on the relationship between the liquid film width produced at the bonding temperature, particle diameter and inter-particle spacing. When the liquid film width is large enough that sufficient particulate material is contained in the melt, particles will be pushed ahead of the solidifying liquid-solid interface resulting in particle segregation at the bond-line. However if the liquid film width is less than some critical value, segregation should not occur. WDS analysis across the joint zone as a function of bonding time indicated that the Ni volume at the joint center varied between 3.67 wt% after 1 minute and 0.15 vol.% after 30 minutes bonding time. The relationship between the width of segregated region and the maximum liquid width that formed during bonding was determined by using equation 12.

$$\mathcal{S}\_{sz} = \mathcal{W}\_{\text{max}} \left[ \frac{\delta\_p + \mathcal{X}\_2}{\delta\_p + \mathcal{X}\_1} \right] \tag{12}$$

Where δp is the average particle size, 1 is the inter-particle spacing in the as received MMC and 2 is the inter-particle spacing after bonding at the joint. By substituting δp = 28 μm, 1 10*m* and <sup>2</sup> 0 into Equation 6.3, the relationship between the width of the particle segregated zone and the maximum width of the eutectic liquid phase was found to be, max 0.74 . *sz S W* This means that the width of the segregated zone is approximately 74% of the width of the maximum width of the eutectic liquid phase formed during bonding.
