Microstructure and Mechanical Properties of Nickel-Aluminum Bronze Coating on 17-4PH Stainless Steel by Laser Cladding

17-4PH stainless steel was chosen as the substrate material. The substrate prepared in size of 100 mm × 100 mm × 20 mm was solid solution treated. It was mechanically ground and cleaned with acetone and anhydrous ethanol. Spherical NAB powder particles with the diameter of 45‒100 μm were selected in this experiment to ensure the uniformity and synchronicity of powder feeding. The morphology of NAB powder is shown in Figure 1. The powder was dried at 60 ℃ for 2 h. The chemical compositions of substrate and powder are listed in Table 1.

The LDD system was made up of an optical fiber laser (TRUMPF TruDiode 3006) in a continuous mode with a maximum power of 3000 W and the spot diameter of 3.4 mm, a NC precision machine, a powder feeder and a gas protection device (Figure 2). During executing process, laser beam, powder feeder and NC program were started up simultaneously, and the molten pool was protected from oxidation and pollution by a special nozzle blowing Ar gas to the melt surface directly at a rate of 10 L/min.

To study the influence of different process parameters on depositing quality, laser power was specified as processing variable, while the scanning speed and hatch distance were set as 1.2 m/min and 1.7 mm, respectively, and the laser repetition width rate was 50%. The power density transferred to energy density expression is as follows: where P is laser power, v is scanning speed, h is hatch distance and t is the thickness of the powder. The laser cladding parameters for each sample are presented in Table 2.

A penetration test was taken to detect the surface defects. The penetrant was sprayed on the cladding surface and then penetrated into the defects through capillary force. The developer was sprayed on the surface after cleaning, and the penetrant at the defect would be sucked out and extended to the surface if flaws existed.

To examine the microstructure of the NAB coating, samples were cut into cuboid with the size of 10 mm × 10 mm × 7 mm. Metallographic preparation on cross-section was as follows: grinding on 400, 800, 1200, 1500 and 2000 # SiC paper in turn, polishing with diamond suspensions (3 μm and 1 μm), and etched by a metallographic etchant composed of 5 g FeCl₃ +20 mL HCl+100 mL H₂O. Leica DM1 5000M optical microscope (OM) was used to observe the microstructure of the cladding cross section. XRD analysis was performed on a Bruker D8 Advance X-ray diffractometer to determine the phases in the coatings. The diffractometer worked with Cu K_α radiation (λ = 0.1540598 nm) at 40 kV and 40 mA, and scans were performed over a 2θ range of 10°‒100° with increments of 0.1° and step time of 2 s.

The surface morphology of the cross section was characterized using a TEMSCAN MIRA 3LMH scanning electron microscope (SEM) fitted with an Oxford energy dispersive spectrometer (EDS). Backscattered electron imaging (BSE) was carried out to distinguish the surface topography and surface chemistry.

The hardness was tested at a Vickers hardness tester (Everone MH-5L) equipped with a diamond regular pyramid tip. The distribution of the hardness at the cross section was measured using the load of 200 N and dwell time of 5 s.

The tensile testing samples were cut along the longitudinal direction of the steel- bronze composite with the interface of steel and NAB coating located at the center of the samples. Samples were built with length of 10 mm and thickness of 1.4 mm. The test was performed with CMT 5305GL machine at room temperature with a load speed of 0.1 mm/min.

The XRD profile of the NAB coatings shows very different peaks from the raw powder (Figure 3). The phases presented in the NAB coatings are α-Cu, which could be a solid solution of Cu, Ni, Al, Fe and Cr, and κ phases (intermetallic compounds such as FeAl, Fe3Al and AlNi), as well as Fe phase. However, in contrast to the data of JCPDS (Joint Committee Powder Diffraction Standards), almost all the diffraction peaks have shifted more than 0.5° from their inherent angular positions. The first two diffractions of Cu are 43.41 and 50.56 in 65-9743# card, respectively, but 42.90 and 49.82, actually.

Figure 4 shows the top morphologies of the coatings before and after the penetrating inspection. This bimetallic structure was composited of two metals with different thermal expansion coefficients, so it’s inclined to form flaws as a result of thermal stress. However, no dyeing trace was detected after the penetrating, which illustrates no major pores or cracks were generated during the cladding process. Almost all composite parts bonded well except S4 (Figure 5). There is a hole defect with the about size of 40 μm.

Figure 6 shows the OM image of the cross section on S1. The coating consisted of three areas: bi-metal transition (zone 1), coarse grain zone combined with dark dendrites (zone 2) and acicular grain zone (zone 3).

Figure 7 shows the typical microstructure of S1. The matrix consisted of coarse grains accompanied with dark dendrites. The chemical compositions of the marked points in Figure 7 are listed in Table 3. It’s obvious that the dendritic positions are rich in Fe and Cr, but poor in Cu and Ni. Combined with the XRD results, the light matrix is Cu-rich solid solution, and the dark dendrites are solidified phases of Fe-rich alloy in the coating (Figure 7a). As shown in Figure 7b, globular κ_II and lamellar κ_III precipitate on the boundary of α or β′ grains, and the other smaller dark particles could be designated as κ_IV according to Murray et al. [20]. Square 1‒3# shows the different morphologies of Fe-rich dendrites.

Figure 8 shows the zone 1 and zone 3 of S1‒S4. It can be seen that the microstructures are compact especially there is a good metallurgical bonding between coating and substrate. Thickness of transition zones are 5‒14 μm. Light coarse grains and dark dendritic structures are found in zone 1. The sizes of light grains are basically the same in each sample. The dark coarse dendrites embedded in the light grains. At the highest energy density of 14.12 J/mm³, coarser grains and less fraction of needle-shaped dark phase (a, b) are formed compared to that of lower energy density at 12.35 J/mm³ (c, d). All light grains are smaller in size and the fraction of secondary Widmanstatten structures are greater at 10.58 J/mm³ (e, f) compared to the former cases. At the lowest energy density, more spiculate microstructures formed in zone 3 even in zone 1 (g, h).

The dendrite sizes range from a few microns to twenty microns (Figure 9). There are three types of spherical particles in the matrix, with diameters approximate of 3 μm, 1 μm and 0.1 μm, respectively, which are pointed out by the arrows. The spherical particles randomly distribute in the matrix, especially, some light finer spherical structures precipitate in the former larger dark ones that can be seen in 1#.

Figure 10 shows the hardness values of different samples at different positions. The result shows a general decreasing of hardness from the substrate to the coating and a sudden drop occurred at the transition zone. The hardness values of the coatings vary by a very small amount, which fluctuates between 204 HV_0.2 and 266 HV_0.2. There is no obvious relationship between the hardness and the energy input. Hardness depends on several factors such as grain size, phase type and phase content ratio. There were some abnormal values in the hardness curve, which were most possibly related to the uneven distribution of κ phases. When the κ phases are dense, the hardness is relatively high, and the hardness is relatively low where less κ phases are formed.

Figure 11 plotted the tensile stress-strain curves for S5 and S6. S5 is found to exhibit considerably longer elongation than S6. The ultimate tensile strength of S5 reached 767.80 MPa, and the specimen experienced ductile fracture. In light of the fracture occurred on the NAB body, it illustrated the interfacial strength was higher than that of as-deposited NAB. Micrographs of tensile fracture surfaces of composite structure are shown in Figure 12. Overall morphology of S5 seems cup like depression and a mass of dimples appears due to ductile fracture of the sample (Figure 12a, b). However, fracture surfaces of S6 show debonding at overlap on layer boundaries (Figure 12c, d).

The dynamic of the molten pool under laser irradiation can be regarded as a flow process dominated by gravity, surface tension and capillary force. There is a significant competitive relationship between melt droplet wetting (or spreading) and solidification process. As shown in Figure 5, unfused defects only occurred in S4. The lower laser energy density results in a lower molten temperature and stored energy. The surface tension of the melt at low temperature is higher, and the melt has poorer viscosity, spreadability and wettability. The low energy storage results in a faster solidification rate, that leading the melt droplets to solidify before they are fully spread. The melt droplet tends to form sphere spontaneously, causing balling on the laser scanning track, which generates the unfused defects in the local area of adjacent molten tracks [22, 23].

Considerable shifts of most of the peaks reveal a rich solid solution relationship among Cu, Ni, Fe and Al elements. Such as an infinite solution between Ni and Cu, and a solid solubility of 9% for Al in Cu. According to Bragg’s law (Eq. (2)), the degree of solid solution will cause the peak position (θ) shifts.

When a solid solution is formed, the dissimilar radius of the atom dissolved into the matrix, that can lead to the lattice distortion and enlarge the atomic radius of the solution. If the crystal plane spacing d increases, the θ value will decrease at a certain λ, showing the diffraction peak shifted to the left.

Table 4 lists the chemical compositions of six types of NAB alloys. Compared to UNS C95500 alloy [22] and the UNS 95800 specimen cut from propeller blades [27], the hardness of the laser coatings is higher than NAB materials processed by the traditional method, but some abnormal values exist on points in Figure 10.

The mass transfer and composition mixing of the powder and substrate in the molten pool have an important effect on the hardness. In general, the hardness of α phase is low, that of β' phase is higher, and that of κ phase is the highest which concludes a series of Al(Fe, Ni) intermetallic compounds. There are three strengthening mechanisms that are solid solution strengthening, precipitation strengthening and dispersion strengthening, and one mechanism of weakening hardness, which is related closely to the hardness of copper alloy cladding layer. The solid solution elements such as Al and Ni are present in α phases, and these solute atoms hinder dislocation motion, that solid solution strengthening works. In addition, many spherical particles that without enough time to merge were acted as reinforcement phase embed in the matrix and prevent dislocation migration, resulting in the increase of hardness of the coating. Al, Ni and Fe form the κ phases with high hardness, which is dispersed in the coating matrix and make precipitation strengthening of the second phase [28, 29]. Appropriate amount of Fe and Ni added to the aluminum bronze alloy can refine the grain, and it is fine grain strengthening mechanism of improving the strength and hardness. Due to the dilution of the substrate, a large amount of Fe-Cr melt mixed into the molten pool and solidified into dendritic structures. As such points 1-1 and 2-2, the laser energy applied at these areas happens to be the higher power parameters, which led to a higher degree of dilution and introduces more Fe-rich melt into the coating. The iron dendrite phases were dense and developed in these regions, resulting in a low local hardness.

As can be seen, S5 has a good tensile property. However, the flat regions on S6 revealed the origin location of fracture. There is a possibility of unmelted powder on this laser parameter. Further, the next layer is deposited on the bumpy surface where there were low-lying areas left by the overlap that results in positive defocus of the laser. This situation provided insufficient laser energy density on powder and would increase the range of uncompleted melt region.