3.1. Effect of Bonding Time on the Bond Interface
The influence of bonding time on the microstructure formed within the joint region during diffusion bonding was assessed by varying the bonding time from 15 min to 60 min.
Figure 3A shows the SEM micrograph of a sample bonded at 850 °C for 30 min. The joint appears to represent a homogenous bond between the Ti–6Al–4V alloy and SDSS. However, a thick reaction layer was observed at the interface on the Ti side of the bond. Al
2O
3 particles appear to have formed clusters and segregated to the grain boundary regions within the Ti side of the bond. The highlighted region presented in
Figure 3A shows an additional reaction layer at the SDSS interface. When the bonding time was increased to 45 min, the width of the interface also increased with the formation of two additional phases within the reaction layer containing dispersed Al
2O
3 particles. Microcracking was also observed in this layer of the interface. The fracture may have occurred during the cutting of the sample; these crack patterns, however, suggest that the reaction layer may be brittle. Further increase of the bonding time to 60 min resulted in the formation of four distinct reaction regions at the interface, primarily on the Ti side. While interdiffusion led to the formation of a wide region, labelled A, the nanoparticles appear to have segregated to form a thick layer at the SDSS interface. EDS analysis of the dark grey particles showed that the Al
2O
3 particles reacted with Cu and Fe [
15,
16].
With increasing bonding time, Kirkendall effects were observed to play an important role in the homogeneity of the interfaces. As identified in
Figure 3, the formation of voids occurred at the Cu/Al
2O
3/SDSS interface, and the volume of voids present in this area increased as the bonding time increased. The formation of these voids was attributed to the interdiffusion of Cu and Fe near the bottom of the TiCu
2 layer. The presence of the voids formed at the SDSS interface was attributed to the differences in the rate of diffusion between Cu in Fe and Fe in Cu. Continued interdiffusion during the diffusion bonding process is believed to result in an accumulation of voids at the SDSS interface and blocks the diffusion of Fe, causing the growth of the TiCu
2 layer.
3.2. Effect of Bonding Temperature on the Interface
The effect of bonding temperature on the microstructural evolution within the joint region was evaluated at bonding temperatures of 800 °C, 850 °C, and 900 °C for 30 min bonding time. Samples bonded at temperatures below 800 °C failed during preparation for testing. For a bonding temperature of 800 °C, three distinct reaction layers were observed at the interface as shown in
Figure 4A. These regions were identified by the differences in the shades of the layers formed. The micrograph also shows evidence of incomplete bonding at the SDSS/Cu/Al
2O
3 interface due to the presence of several cavities/voids. On the other hand, the Ti/Cu/Al
2O
3 interface showed evidence of uniform bonding between the Cu/Al
2O
3 interlayer and the Ti base metal. This dissimilarity may be attributed to the differences in the diffusion coefficient of Cu in Fe and Cu in Ti. The literature shows that within the temperature range studied, the diffusivity of Cu in Ti shows fast impurity diffusion when compared to other metals [
17]. Further increase of the bonding temperature to 850 °C led to the formation of a uniform bond at both the SDSS/Cu/Al
2O
3 and Ti/Cu/Al
2O
3 interfaces, as shown in
Figure 4B. The Al
2O
3 nanoparticles appear to be distributed along the grain boundaries on the Ti side of the bond. When the bonding temperature was further increased to 900 °C, the width of the reaction zones at the Ti/interlayer interface increased in thickness, and a third reaction layer was observed at the SDSS/interlayer interface (see
Figure 4C). According to the Ti–Cu phase diagram, bonding within the temperature range of 870–900 °C would allow for the formation of a eutectic liquid of the form L
. Evidence of the formation of the eutectic liquid is shown in
Figure 5A. Additionally, at higher bonding temperatures the
microstructure forms at the Ti/Cu/Al
2O
3 interface (see
Figure 5B)
3.3. Effect of Interlayer Composition on the Bond Interface
The influence of the interlayer composition on the microstructure of the joint region was investigated by comparing bonds formed using an electrodeposited Cu coating containing α-Al
2O
3 nanoparticles (
Figure 6A,B) and a 25 µm Cu foil was used as the interlayer (
Figure 6C,D).
Figure 6A shows the presence of six separate distinct layers, where each layer is distinguishable by the shade of the specific region. These areas were labelled P
2–P
7. The elemental composition of each phase was identified using EDS and is presented in
Table 2. The EDS analysis shows that the composition of P
1 is that of the super-duplex stainless steel because of the high Fe and Cr content and low Cu concentration.
The dark grey region P
2 is a Cu-rich phase of TiCu
2 which contains 67 wt % Cu and 28 wt % Ti. P
3 also appears to be a Cu-rich phase with a composition of 33 wt % Ti and 58 wt % Cu and is possibly the same TiCu
2 (see
Figure 6B). Similar changes were observed for region P
4, which had a composition of 41 wt % Ti and 55 wt % Cu, while P
5 similarly showed a marginal change in the Ti content. The EDS analysis suggests that both P
4 and P
5 are TiCu compounds (see
Figure 6A). The Ti-rich phase P
6 was identified as Ti
2Cu into which Cu had diffused. The general trend showed that the Ti content within the reaction layers increased progressively from the steel side of the bond towards the Ti base metal. The region labelled P
7 contained spherical particles distributed within the grain boundary region. EDS analysis showed the presence of Al
2O
3 particles as well as Cu, Ti, and Fe.
The Ti/SDSS joint bonded using 25 µm Cu foil presented in
Figure 6C shows the presence of five distinct reaction regions identified by the differences in shade, similar to those identified when the Cu/Al
2O
3 interface was used. EDS analysis of the region labelled P
8 suggests that this phase is a ternary compound of the form Ti
xCu
x-Fe
x. The region labelled P
9 is believed to be compound TiCu given the high Ti content recorded (see
Figure 6D). A summary of the phases identified is listed in
Table 2. The low 0.54–1.95 wt % of Fe recorded within the reaction layer confirms that phases P
3, P
4, P
5, and P
7 are likely binary compounds formed from reactions between Ti and Cu. EDS maps of the bond formed using Cu foil as the interlayer showed that the distribution of the elements Fe, Cu, and Ti within the reaction layer (
Figure 7) led to the formation of the ternary intermetallic compound labelled at P
8: (Ti
33Cu
67 −
xFe
x; 1 <
x < 2.5), T
2 (Ti
40Cu
60 −
xFe
x; 5 <
x < 17), T3 (Ti
43Cu
57 −
xFe
x; 21 <
x < 24), τ (Ti
37Cu
63 −
xFe
x; 6 <
x < 7), and
τ1 (Ti
45Cu
55 −
xFe
x; 4 <
x < 5) [
18].
EDS line scans presented in
Figure 8 and
Figure 9 show a visual representation of the variation of the composition across the joint zone and confirm the high concentration of Ti and Cu within the reaction layers. The Ti–Cu phase diagram shown in
Figure 10A suggests that within the range of composition listed in
Table 2, the compound TiCu
2 would form with Cu content varying between 33 wt % and 48 wt %. The compound labelled P
8 was only found at the interface when Cu foil was used as the interlayer. This phase has a higher concentration of Fe and is believed to be a ternary intermetallic phase.
These phases are believed to have formed due to the dissolution of approximately 38 wt % of Cu in the cubic Ti–Fe lattice. The Ti–Cu–Fe ternary phase diagram presented in
Figure 10B confirms the likelihood of the formation of intermetallic compounds within the temperature range of 800–900 °C [
19].
The shape of the compound formed at the interface when the Cu foil was used as the interlayer appears to be a continuous series of brittle plates, which are stacked as shown in
Figure 6D. When compared to the bond formed using the Cu/Al
2O
3 coating as the interlayer, the nanoparticles are seen to occupy positions around grain boundaries as shown in
Figure 6C.
The EDS line scans presented in
Figure 9 show that a large portion of the reaction layer formed on the Ti side of the bond was promoted by the diffusion of Cu into the Ti base metal, given that Cu diffuses faster into Ti than it does into SDSS. Similar findings were observed when the Cu/Al
2O
3 interlayer was used (see
Figure 9). The reaction layer was found to have a higher percentage of Cr and Fe, which may suggest that Fe and Cr diffused faster into the Cu/Al
2O
3 coating than into the Cu foil.
When the process was modelled using Thermo-Cal, the results showed that within the bonding temperature range studied (850–900 °C) four binary phases are possible; Cu3Ti2, Cu
2Ti, Ti
2Cu, and Cu
4Ti
3, as shown in
Figure 10A, support the EDS data. Similar observations were reported by Pardal et al. [
3] who studied dissimilar metal joining of Ti and duplex stainless steel using Cu as a transition metal interlayer. A plot of the Gibbs free energy profile shown in
Figure 10C,D indicates that several stable phases are likely to form in the Ti–Cu binary system for the bonding temperature range studied. The sequence of phase formation goes from CuTi
2 to Cu
4Ti
3, Cu
3Ti
2, and Cu
2Ti. The growth of these phases is believed to be promoted by interdiffusion between Cu and Ti into the reaction layer. The XRD spectrum presented in
Figure 11 confirms the presence of Cu
2Ti within the bonded region.
3.4. Hardness
The influence of bonding temperature on variation of the hardness across the bonding interface was evaluated and is presented in
Figure 12. The hardness values were measured across the joint starting at 300 µm from the joint center. The results show that the hardness of the Ti base metal fluctuated between 320 and 420 VHN up to 100 µm from the joint center as the bonding temperature was increased from 800 to 900 °C. The hardness within the joint center was also observed to decrease from 634 to 405 VHN when the bonding temperature was increased from 800 to 900 °C. The hardness of the SDSS base metal was found to be higher than that of the Ti base metal, with a hardness value of 387 VHN. The result of the hardness test shows (see
Figure 12A) that the hardness within the joint center decreased with increasing bonding temperature. These findings were attributed to increased diffusion rates at higher bonding temperatures leading to homogenization of the bond region.
The effect of bonding time on variation of the hardness across the joint interface was also studied (see
Figure 12B). Similarly, the hardness was measured across the joint starting at 300 µm from the joint center as a function of bonding time. The hardness within the joint center increased from 554 to 752 VHN when the bonding time was increased from 30 min of bonding time to 60 min. The maximum hardness of 750 VHN was recorded at the center of the joints bonded for 60 min. The hardness at the center of the bond is believed to have been caused by the formation of the reaction layer made up of binary and ternary intermetallic compounds as well as Al
2O
3 nanoparticles dispersed at the joint interface [
20].
Finally, the impact of interlayer composition on variation of the hardness measurements (see
Figure 13) across the joint region was evaluated by comparing the bonds made with Cu/Al
2O
3 and Cu foil interlayers. The results show that when the Cu/Al
2O
3 interlayer was used, higher hardness values were recorded in the joint region when compared to samples bonded using Cu foil. The results show that the hardness within the joint center was marginally higher when Cu/Al
2O
3 coating was used as the interlayer due to the dispersion of hard nanoparticles at the interface.
3.5. Growth Kinetics of Interfacial Phases
The width of the reaction layer was measured as a function of bonding time and bonding temperature and plotted as shown in
Figure 14. The data indicate that as the bonding time increased, the width of the reaction layer also increased as predicted by the power law shown in Equation (1). Similar behavior was observed when the bonding temperature was increased from 800 to 900 °C. This response is believed to be consistent with the changes in the rate constant as the bonding temperature increased. The width of the reaction layer increased with increasing bonding temperature and can be estimated by the mathematical relationship shown in Equation (1):
where
d is the thickness of the reaction layer,
k is the rate factor,
t the diffusion time, and
n the time exponent. Further,
, where T is the bonding temperature (K), Q is activation energy (kJ/mol), R is the real gas constant (8.314 J/K mol), and
ko is the growth constant (m
2/s). The growth kinetics of the intermetallic layer is controlled by interdiffusion (volume diffusion); therefore, the diffusion time is estimated to be
t1/2 (i.e.,
n = 0.5). On the other hand, if the growth kinetics was controlled by interfacial diffusion, the time exponent would be
n = 1 [
21,
22].
The results of the experimental study were used to assess the kinetics leading to the formation of the Ti
2Cu reaction layer at the interface. A summary of the growth kinetics of the reaction layer is shown in
Table 3. The activation energy was calculated to be 111.4 kJ/mol for a bonding temperature of 850 °C. This is similar to findings published in the scientific literature which suggest that the activation energy for the formation of Ti
2Cu is approximately 122.1 kJ/mol [
16]. In this study, the calculation of the activation energy was performed based on a simplified diffusion model, which assumes diffusion in a single phase which is influenced by the thickness of the reaction layer formed during the diffusion bonding process and the increases in the rate coefficient. The width of the diffusion layer is also believed to be driven by the diffusion of Cu into the Ti base metal, leading to the formation of various compounds. The Thermo-Cal model predicted the possible formation of several other compounds at the interface; however, the limitation of the simplified diffusion model is that only a single phase is assumed to have formed at the interface.
3.6. Mechanism of Bond Formation
The diffusion bonding process involves several stages [
13]. The first stage involves initial contact between the faying surfaces and the interlayer, followed by deformation of the surface asperities under the effects of an externally applied load. The combination of heat and load increases the area of contact between the base metals and the interlayer and closes some of the voids or cavities present at the interlayer/base metal interfaces.
The second stage of bonding occurs due to interdiffusion between the base metals and the interlayer, leading to the formation of an intermetallic compound at the interface, which promotes the closing of remaining voids and cavities. The third stage of bonding involves diffusion of Cu into the SDSS and Ti–6Al–4V base metals, leading to a eutectoid reaction as predicted by the equation
for bonding temperatures of 850 °C and below. The EDS analysis presented in
Table 2 shows that intermetallic compounds such as CuTi
2 and Cu
2Ti are also formed at the interface during the bonding process. A schematic of the mechanism involved in the bond formation is shown in
Figure 15. The results show that the width of the reaction layer increases with increasing bonding time. Similarly, the literature shows that the strength of the bond increases with increasing thickness of the reaction layer. When a bonding temperature of 900 °C is used, a eutectic liquid forms at the interface, as illustrated in
Figure 15E, and solidifies isothermally in the subsequent stages of the bonding process. At temperatures below 900 °C, solid-state bonding is achieved.