Microstructural analysis, micro-hardness and wear resistance properties of quasicrystalline Al–Cu–Fe coatings on Ti-6Al-4V alloy

The substrate material used in the present investigation was Ti-6Al-4V alloy with the chemical composition (wt%) 6.2 Al, 3.9 V, 0.2 Fe, 0.12 C, 0.15 O, 0.005 N, 0.0003 H and Balanced Ti. The substrate was cut and machined into dimensions 70 × 70 × 4 mm³. Prior to laser treatment, the substrates were sandblasted, washed, rinsed in water, cleaned with acetone and dried at room temperature before exposure to laser beam to minimize reflection of radiation during laser processing and enhance the absorption of the laser beam radiation. Al (99.8% purity), Fe (99.97% purity), Cu (99.9% purity), reinforcement metallic powders were used as alloying powders mixed in Al-3Cu-5Fe (A₁), Al-3Cu-7Fe (A₂), and Al-3Cu-10Fe (A₃) ratios, respectively, in a shaker mixer (Turbular T2F; Glenn Mills, Inc.) for 18 h at a speed of 49 rpm to obtain homogeneous mixture. The particle shape of the powder used was spherical with 50–105 μm particle sizes. Samples were characterized for SEM and energy dispersive spectroscopy (EDS) analysis. Specimens for SEM (JSM-7600F; JOEL, Ltd) were prepared by cutting samples in such a way to reveal the transverse section of the coatings. While the microhardness was measured using four indentations and the average microhardness was recorded (figure 1).

Zoom In Zoom Out Reset image size

Laser surface alloying was performed using a 3 kW continuous wave (CW) Ytterbium Laser System (YLS) controlled by a KUKA robot which controls the movement of the nozzle head and emitting a Gaussian beam at 1064 nm. The nozzle was fixed at 4 mm from the steel substrate. The admixed powders were fed coaxially by employing a commercial powder feeder instrument equipped with a flow balance to control the powder feed rate. The metallic powder was fed through the off-axes nozzle fitted onto the Ytterbium fibre laser and it was injected simultaneously into a melt pool formed during scanning of the Ti-6Al-4V alloy by the laser beam. An argon gas flowing at a rate of 2.0 L min⁻¹ was used as a shielding gas to prevent oxidation of the sample during laser surface alloying. Overlapping tracks were obtained by overlapping of melt tracks at 70%. To determine the best processing parameters, optimization tests were performed with the laser power of 800 to 900 W and scanning speed varied from 0.6 to 1.2 m min⁻¹. The final selection criteria during optimization tests was based on surface having homogeneous layer free of porosity and cracks determined from SEM analysis. The optimum laser parameters used was 800 and 900 W power, a beam diameter of 4 mm, gas flow rate of 2.0 L min⁻¹, powder flow rate of 2.0 g min⁻¹ and scanning speeds of 1.0 and 1.2 m min⁻¹ respectively.

Figure 2 illustrates the grain structure of titanium alloy. The average grain size is approximately 76 μm. The microstructure of substrate is represented by light and dark regions indicating two phases namely the alpha (α) and beta (β) phase that are equally distributed through the area of the substrate as shown in figure 2. This is the basic structure of titanium substrate that is used in this investigation.

Zoom In Zoom Out Reset image size

The high-resolution microstructure is achieved through SEM. Figure 3 illustrates the high magnification of composites formation and the bonding between the coating and the substrate at laser power of 800 W and scanning speed of 1.0 m min⁻¹. Various phases that are differentiated by their colour and shape can be observed. Analysing figure 3, the lath-shaped Al₃Fe monoclinic λ−phase can be found throughout the clad. The solidification mechanism is due to laser beam interaction, the coating powder particles melt and create a liquid pool on the titanium (substrate) surface that undergoes a limited melting [22]. Therefore, only small dilution of addition elements is obtained. Solidification begin with formation of some large particles such as Al and Fe. Since Cu is well known as a strong β-stabilizing element, it can be observed that atomic migration of Cu into Ti lattice results in the β-Ti phase formation during cooling and travels a longer distance in the Ti lattice than other elements which opens more crystallographic structure of the β matrix [22]. Aluminium and copper alloys are not really perfect materials for process of LMD due to lower absorption of photons and high reflectivity as compared to LMD friendly materials such Fe based alloys. Thus, aluminium and copper powder are coated with Fe for better photons absorption resulting in better deposit quality. This alternate layer morphology demonstrates the presence of eutectic phase in the inter-dendritic regions. It is found that low solidification velocity and high thermal gradient at the bottom of melt pool and above the solid/liquid interface results in the presence of columnar dendrites. An equiaxed microstructure is developed from middle to top section of the composite and this due to high solidification velocity and low thermal gradient at the upper part of the melt pool. This is well known phenomenon in rapid solidification [23].

Zoom In Zoom Out Reset image size

Figures 4 and 5 represent the microstructure of Ti6Al4V/Al-Cu-10Fe and Ti-6Al-4V/Al-Cu-7Fe composites. It illustrates the microstructures of Ti6Al4V/Al-Cu-10Fe composites at laser power of 800 and 900 W respectively. In this region, where Fe concentration is highest, the incident heat and photons from melt pool initially interact with the Fe coating and absorbed heat energy conducted through the core and gradually interacts with titanium and aluminium. The microstructure shows a classic ferrite equiaxed grain and it is elongated nearly along the direction of cooling due to the high rate of cooling in rapid solidification [23–26]

Zoom In Zoom Out Reset image size

Zoom In Zoom Out Reset image size

It can be observed that laser power has significant influence on microstructure of these composites. Increasing laser power increases the columnar grains as well as the dendritic microstructure. At lower laser scanning speed under constant laser power, the microstructure is coarse but by increasing laser scanning speed, the microstructure is refined. However, at very high laser scanning speed, it results in cracking susceptibility and an increase in temperature gradient of interface such as figure 4(b). These cracks propagate through the grains and on the grain boundaries and have no regularity. They initiate on the solidification stage in the HAZ (heat affected zone) of adjacent tracks (under layers) or on the grain boundaries indicating that there are high stresses in the cooling and solidification zones as well as in the HAZ zone which provides propagation and origination of cracks. It is evident that an increase in laser power results in an increase in degree of melting therefore, decreasing low fusion porosity indicating that the quantity and size of pores were reduced. The coating is directly deposited on high thermal conductive Ti6Al4V substrate solidifying the melt very rapidly with significantly high cooling rate and undercooling thus, few pores were formed.

EDS was utilised to analyse the element concentration distribution along the transition composition route. Figure 6 illustrates the number of elements at different scanning speed and laser powers. It can be observed that there was higher number of titanium and aluminium presented as per the theoretical expectation. The diffusion of all these elements are due to high-temperature gradients and multiple heating in laser metal deposition.

Zoom In Zoom Out Reset image size

There is a proportional relationship between laser power and geometrical properties of the clad. Increasing laser power will increase the clad and HAZ height of the coatings. This was expected in accordance with the literature review; an increase in laser power increases the clad height. This takes place due to the laser-material interaction during the laser metal deposition. The un-melted powder particles initially travel through the laser beam before landing into the melt zone. This causes powder to melt before reaching its melt-pool however it reduces some of the laser power, which means less power is available to melt the substrate. This causes a narrower melt pool to be created indicating that the width of the clad would also be narrower. This can be seen by analysing the graph of deposit height and width vs laser power (figure 7). At laser power of 800 W, the width is smaller but by increasing laser power to 900 W, there will be a direct increase in the energy input level that allows for more power to be available for the melt pool to be created and therefore width becomes wider. The height of the clad however continue to increase due to sufficient energy to melt the incoming powder particles on the substrate. Hence, it can be understood that the influence of laser power on the clad width is comparatively more significant than its influence on the clad height [27, 28]

Zoom In Zoom Out Reset image size

Plotting the graph of deposit height and deposit width as a function of scanning speed (figure 8), it can be seen that these geometrical properties decrease with increasing scanning speed. This is due to a decrease in the amount of powder delivered per unit length, a reduction in the interaction time between the laser and the substrate as well as the energy input per unit length. It is evident that an increase in laser scanning speed results in an increase in aspect ratio. This is due to the fact that increasing laser scanning speed decreases the clad height while it has a slight effect on the clad width. Hence, it can be understood that the influence of laser scanning speed on the clad height is comparatively more significant than its influence on the clad width. It can be concluded that increasing laser scanning speed decreases the efficiency of powder. This is due to the width of melt pool becoming narrower as the powder efficiency decreases and laser scanning speed increases. Therefore, it reduces the possibility of powder catchment as a result of an increase in less favourable impact conditions (Liquid particle-solid substrate) or un-favourable impact conditions (Solid particle-solid substrate) [29].

Zoom In Zoom Out Reset image size

The morphologies of the laser cladded samples (figure 4) displayed different characteristics as a result of varied laser power and different cooling rate. Also, there is a difference in the heat affected zone (HAZ) bands. Observation shows a decrease in the grain size of the microstructure with the increase in processing scan speed within the coatings. As regards to solidification and formation of dendrites as evident in figures 4(a) and (b), when the laser scanning speed is low (1.0 m min⁻¹), samples engage with the absorbed heat for a longer period of time and the cooling rate is slower, therefore grain growth takes a longer time to its own accord and consequently resulting in coarse microstructures and larger band of the HAZ [30]. However, at higher scanning speed (1.2 m min⁻¹), the interaction time with the absorbing energy of the substrate are reduced. Consequently, this result in high cooling rate and a corresponding rapid nucleation rate yielding smaller sizes of grains after solidification. Simultaneously, The HAZ of high scanning speed samples showed a thinner band as compared to low scanning speed samples. The results can be related to the works of Qin et al [31] and Chen et al [32]. Thus, it can be agreed that the grain sizes of the cladded samples can be refined by increasing the scanning speed. The dendritic and inter-dendritic structures are also observed from the matrix. At high laser scanning speed during cladding, the cooling rate of the melt pool is very high, resulting in larger dendrite arm spacing as growth is observed towards the substrate (figure 4). Thus, it is well observed that the fast cooling rate resulting in fine dendritic microstructure and low dilution increased the micro-hardness properties of the coating [33]. Conversely, at lower scan speed (1.0 m min⁻¹) coarse microstructure with longer interaction time is exhibited. Furthermore, it was also observed at figure 4 that laser induced thermal stress occurred at the transition zone leading to formation of cracks.

The XRD results for Ti6Al4V/Al-Cu-5Fe and Ti6Al4V/Al-Cu-10Fe composites at different laser power and scanning speeds are shown in figures 9 and 10 respectively. Analysing Ti6Al4V/Al-Cu-5Fe graph, it was found that due to the effect of dilution, a part of Ti entered into the molten pool from the substrate. The results indicate that Ti, Al₃Ti, Ti₃Al, CuTi₂ is formed through the in situ metallurgical reactions during the laser metal deposition process. This was expected in accordance with the results achieved from EDS (figure 6) and SEM (figure 3) analyses. It can be observed that no Fe seems to be presented as laser power and scanning speed increase to 900 W and 1.2 m min⁻¹ respectively. Increasing laser power and scanning speed will significantly increase the diffraction peak of Ti and Ti₃Al. Since there is no indication of Fe or Cu, Al and Ti concentration is highest which leads to the formation of titanium aluminide (Ti₃Al) and presence of α & γ phases. Evaluating Ti6Al4V/Al-Cu-10Fe graph (figure 10), it is evident that the elemental Ti peaks were not observed. This indicates that Ti reacted with increased amount of Al and Fe and less amount of Cu powders to form Al₂Ti, Fe_0.975Ti, Al_0.9Fe_3.1Ti_0.025 and AlCu₂Ti. Increase in amount of Al and Fe and decrease in amount of Cu results in a change from fcc structure to a combined fcc and bcc structure. Thus, Al and Fe promotes bcc phase formation while Cu promotes fcc phase formation. Aluminium has a larger impact on causing the change in this structure.

Zoom In Zoom Out Reset image size

Zoom In Zoom Out Reset image size

Crystalline phase with dissolved copper were found in the XRD. Quasicrystalline phases were also formed. The crystallization of the quasicrystalline phases occur during isothermal heat treatment. Rapidly solidified alloys are mainly composed of a solid solution of alloying elements and stable phases whose presence depends on the content of alloying elements. Rapid solidification leads to the formation of stable intermetallics and quasicrystalline phases. Attractive mechanical and wear properties are the attributes of quasicrystalline Al–Cu–Fe alloys. These properties can offer huge potentials of industrial applications. At the interfacial zone between the substrate and the coatings, partial melting occurred, and the heat slightly affect the substrate. There was minimal dilution due to existence of melting of thin zone at the substrate. Despite the minimal dilution and melting, there was metallurgical bonding formation between the substrate and the coating. The enhancement in hardness had been reported to be due to solid solution strengthening as a result of titanium concentration in copper. The intermediate and outermost layers consist of TiCu and Ti₂Cu intermetallic phases. Ti₃Al has also been reported to have improved wear resistance properties [34].

The hardness values were recorded at different points on the intermetallic section where formation of the composite takes place. Then the average of these values for each sample were recorded. Figure 11 illustrates the microhardness of Ti6Al4V/Al–Cu–Fe coatings. At scanning speeds of 1.0 m min⁻¹ and 1.2 m min⁻¹, it is evident that the mean hardness value decreases with increasing laser speed. Ti6Al4V/Al-Cu-5Fe compared to other composites has the highest mean hardness value at laser power of 800 W. Higher percentage of Al and Cu and less amount of Fe can increase the hardness at lower laser power. It was found that the lowest hardness value appears in the region of Ti6Al4V substrate and as moving along the direction of deposition, the hardness value begin to increase significantly. This non-linear increase is due to the diffusion of the coating and discrepancy between unmelted and resolidified coating hardness. It is seen that by using high laser power cooling rate decreases, thereby forming coarse dendrites that reduces the hardness.

Zoom In Zoom Out Reset image size

However, this is not true for Ti6Al4V/Al-Cu-7Fe and Ti6Al4V/Al-Cu-10Fe composites (figure 12). It can be seen that for both scanning speed of 1.0 m min⁻¹ and 1.2 m min⁻¹, the mean hardness value increases with increasing laser power. In this case, at the start of material deposition, the heat rapidly dissipated by the substrate heat conduction. At the beginning stage of the LMD process, the initial thermal transience produced a rapid quenching rate effect which resulted in hardness to increase. It can be observed that the mean hardness value increases with increasing laser power and this is due to the limited interaction within the powder and the laser; only a low amount of energy was supplied during the melting process. This low energy limited the amount of time available for growth, thus increasing the hardness. It is evident that both composites have the highest mean hardness values at higher laser power (900 W) and lower scanning speed (1.0 m min⁻¹).

Zoom In Zoom Out Reset image size

Al₂Ti, had been reported to be practically overlooked. Al₂Ti, like Al₃Ti, possessed lower density and better oxidation resistance compared to Ti₃Al and TiAl based alloys as reported in the literature [24, 34]. Al₂Ti produces material with a finer microstructure, a higher hardness, and increased resistance to room temperature cracking as shown in figure 10.

Figures 13 and 14 depict the graphs of coefficient of friction (COF) and wear depth (WD) as a function of time for the substrate and composites respectively. This is accomplished at different laser power and scanning speed. Evaluating the graphs of COF vs time for the composites and substrate, it can be observed that COF suddenly reaches its maximum for all the graphs and as the time begins to increase, the COF values start oscillating within that limits with increase in time. However, this is not true for the graph of wear depth vs time as it is evident that the wear depth increases gradually with increasing time. Ti6Al4V/Al-Cu-5Fe composite has the lowest coefficient of friction. This was expected in accordance with results obtained from hardness testing; high mean hardness value result in low coefficient of friction and low wear rate. This also implies that Ti6Al4V/Al-Cu-5Fe can be used in the aerospace industry for components such as engine turbine blades and bearings. Wear depth begins to increase as there is an increase in Fe concentration (figures 14 and 15). This can be seen from the graphs of wear depth vs time, the wear depth values for both the substrate and Ti6Al4V/Al-Cu-5Fe composite is lower than 0.025 mm and for the Ti6Al4V/Al-Cu-7Fe & Al-Cu-10Fe composites are above 0.03 mm. Figure 15 presents the coefficient of friction (COF) with sliding time of reinforced admixed laser clad Al–Cu–Fe metallic coatings and as-received Ti-6Al-4V alloy. The test was carried out under conditions of dry sliding situation of a load of 15 N at a sliding velocity of 2 m s⁻¹ and 2 mm sliding distance using tungsten carbide ball. The entire samples were sectioned to dimensions 1 cm × 2 cm area that can be fixed securely in a fitted sample chuck. The coefficient of friction was recorded continuously during the test for duration of 1000s (17 min) of reciprocating movement. The pattern implied that the COF of Ti-6Al-4V alloy increased at initiation and later fluctuated.

Zoom In Zoom Out Reset image size

Zoom In Zoom Out Reset image size

Zoom In Zoom Out Reset image size

Friction coefficients were observed to be reduced with the incorporation of the metallic coatings. Al–Cu–Fe ternary coatings exhibit the best wear resistance performance as a result of increased rate of solidification of laser deposited coatings which, in turn, produce matrices with smaller grain size. Consequently, the smaller the grain size the higher the resistance of the coating to deformation, which is according to the Hall-Petch equation [35]. Therefore, the reduced grain size of coatings with higher scan speeds enables the coatings to possess a reduced contact area of the coating surface with the steel ball, thus mitigating the smearing effect of the coatings on the steel ball. Lack of surface coating on the substrate is responsible for the high coefficient of friction recorded for the Ti-6Al-4V substrate.

The friction coefficient is calculated using the applied load and the frictional force, this is automatically done from the instrument by the following equation:

$$\begin{eqnarray}&&\mu =\displaystyle \frac{{F}_{f}}{{F}_{N}}\end{eqnarray} \tag{ 1 }$$

Where: μ = friction coefficient (dimensionless) F_f = Frictional force (N) and F_N = Normal load (N)

Figure 16 shows the wear scar obtained during the wear test for the samples. According to Qu and Truhan [36], the wear depth Zw and wear volume of the flat specimen can be calculated with the following equations:

$$\begin{eqnarray}&&{V}_{W}={L}_{S}\left[{{R}_{S}}^{2}{\sin }^{-1}\left(\displaystyle \frac{W}{2{R}_{S}}\right)-\displaystyle \frac{W}{2}({R}_{S}-{Z}_{W})\right]+\displaystyle \frac{\pi }{3}{{Z}_{W}}^{2}(3{R}_{S}-{Z}_{W})\end{eqnarray} \tag{ 2 }$$

$$\begin{eqnarray}&&{Z}_{W}={R}_{S}-\sqrt{{{R}_{S}}^{2}-\displaystyle \frac{{W}^{2}}{4}}\end{eqnarray} \tag{ 3 }$$

Where Z_W = Wear depth; R_S = Radius of spherical surface at both ends; W = Width of the wear scar V_W = Wear Volume; L_S = Stroke length

Zoom In Zoom Out Reset image size

Graphs of COF as a function of time (figure 15) for all the composites and substrate show a similar trend except at laser power of 900 W, the graphs have the steepest initial gradients and the Ti6Al4V/Al-Cu-10Fe composite has the least steep initial gradient.

The wear track for a Ti6Al4V/Al–Cu–Fe composite is illustrated in figure 16. The average wear track thickness is 410.7 μm. Wear thickness is an essential parameter used in industries to compute the wear rate, enabling to predict the life expectancy of a component that is subjected to wear.

The tribometer was used to measure the wear depth and COF while an optical microscope was utilized to measure the thickness of the wear track as shown in figure 16. The results obtained from wear track thickness is used in order to calculate the wear rate. However, comparing the calculated wear rate values lead to inaccurate results since the wear volume is not identical for all the samples. Unlike COF and wear depth, wear thickness increases with increasing laser power from this study. It can be observed that substrate has the highest wear thickness value compared to the Ti6Al4V/Al-Cu-5Fe, Al-Cu-7Fe and Al-Cu-10Fe. Reduction in wear thickness (figure 17) increases the contact stress. Ti6Al4V/Al-Cu-10Fe shows the highest wear thickness which indicates increase wear loss. However, adding higher amount of Fe into the coating can increase the wear thickness of the composite. Ti-6Al-4V/Al-Cu-10Fe coating displayed the highest wear thickness and it also indicates highest wear volume loss. There was wear reduction of 47.46%, 30.51% and 16.95% in Al-Cu-5Fe, Al-Cu-7Fe and Al-Cu-10Fe respectively. Al-Cu-5Fe coating shows the highest wear resistance performance compared to the substrate. Al-Cu-5Fe coating was 2.8 times the wear resistance performance of the substrate. Hence, the wear resistance performance of Ti-6Al-4V/Al-Cu-5Fe coating which also has the highest hardness is the best among all the ternary coated samples [37, 38].

Zoom In Zoom Out Reset image size

The depth of groove and magnitude of delamination increases as the coating thickness increases, since the hardness of coating is a function of wear resistance. Formation of microcracks were observed in all the coatings but microcracks were more distinct in the highest coating thickness. This is because, the higher coating thickness has higher porosity and low micro-hardness values. The low-microhardness and high porosity in the coating results for the higher wear volume loss [39–46]. Therefore, the wear volume loss increases as the coating thickness increases.