Microstructure Analysis of HPb59-1 Brass Induced by High Current Pulsed Electron Beam

Surface and cross-sectional analysis

Figure 1 shows the morphology of original structure of the HPb59-1 brass after etching. The surface microstructure of the original sample was relatively coarse, and the grain size ranged from several to dozens of microns. At room temperature, it consisted of black strips of α (Cu–Zn solid solution with face-centered cubic space lattice), lightly colored gray β′ (solid solution on the basis of electronic compound CuZn with body-centered cubic space lattice) and scattered white Pb phases. Pb was distributed on the grain boundary and the bases of α and β′ in the form of free states.

Figure 1:

SEM image of the original structure of the HPb59-1 brass after etching.

The surface morphologies of lead brass after HCPEB treatment under different pulses were shown in Figure 2, which presented a corrugated fluctuant characteristic. The surfaces of magnesium alloy [9] and NiTi alloys [10, 11] with the treatment of HCPEB also have analogous morphologies mainly due to the evaporation of alloy elements. Besides, a small number of craters were formed on the surface layer after 5 pulses treatment. The formation mechanism of caters on the metal surfaces induced by HCPEB have been investigated in many previous works, and this phenomenon results from the nonhomogeneous local melting and subsequent eruption occurred in the subsurface layer of the target material after HCPEB treatment [12]. With the increase in number of pulses, fewer craters could be observed and the corrugated fluctuant feature of morphology became more apparent after 25 pulses, which was related to the amount of energy injected. Crater morphology could not be preserved with the continuous evaporation.

Figure 2:

The surface morphology of HPb59-1 copper alloys after HCPEB treatment at (a) 20 kV, 5 pulses; (b) 20 kV, 15 pulses; (c) 20 kV, 25 pulses and (d) microcracks with 20 kV, 15 pulses.

As shown in Figure 2(d), after HCPEB bombardment, microcracks appeared on the alloy surface due to the great thermal stresses induced by HCPEB. The defects of crack in the microstructure made an adverse effect on the surface performance of the alloys, which should be avoided by increasing the pulse width of electron beams [13].

Figure 3 shows morphologies of the cross section of HCPEB-treated HPb59-1 alloys. The remelted layers got thickened as the number of HCPEB pulses increased. For 5 pulses (Figure 3(a)), the thickness of remelted layer was about 3.0 µm, whereas it reached about 4.2 µm after 25 pulses treated (Figure 3(b)). The effect of heat accumulation in underlying substrate materials was considered to be a reasonable explanation for the increasing thickness of the remelted layer with multiple short interval pulses. The supersaturated solid solution formed on the surface, which was conductive to the improvement of corrosion resistance due to the rapid melting and solidification process [9].

Figure 3:

Cross-sectional SEM morphologies of HCPEB-treated HPb59-1 copper alloys after etching (a) 20 kV, 5 pulses and (b) 20 kV, 25 pulses.

Distribution of Pb

Some bright white substances were observed on the alloy surface (Figure 4(a) and (b)), which was in a Pb-rich phase with a small amount of Zn and Cu elements according to the energy-dispersive X-ray spectroscopy (EDS) of the corresponding point 1 (Figure 4(c)). With relatively lower melting points, Pb and Zn on the surface of the copper alloys would have a priority of evaporating due to the rapid heating and melting bombarded by HCPEB. The solid solution of Pb in Cu and Zn was rather low, while Zn in Cu could reach up to 39%. Thereby the vapors of zinc deposited onto the surface as well as the supersaturated solid solution of Cu and Zn was formed. However, Pb was “rejected” owing to the extremely low solubility and stacks onto the blocks under the effect of coagulation, showing a bad associativity with the matrix.

Figure 4:

SEM fissure morphology images of Pb on HPb59-1 alloy surface after HCPEB treatment with 15 pulses, 20 kV (a) low resolution; (b) high resolution; (c) EDS result on the corresponding Pb area and (d) Pb particles.

The morphology of pure Pb particles was observed on the HPb59-1 alloy surface after HCPEB treatment with 15 pulses (Figure 4(d)). At high power, many small round particles were distributed on the surface with the size of submicron levels, which were proved as pure Pb particles by the EDS analysis. The generation of these particles was mainly attributed to the vaporization of Pb on the surface [7].

Nanostructure analysis

Fine particles were distributed massively on the alloy surface, the chemical composition of which was determined as pure Pb through the EDS analysis. After HCPEB treatment with 15 pulses, many nanosized Pb with a general grain size between 50 and 100 nm were observed (Figure 5(a) and (b)). Pure Pb was vaporized to the treated surface layer. The formation of this morphology feature could be explained as follows: when the HCPEB-treated surface temperature reached the Pb boiling point, vaporization occurred at a given depth determined by the energy density of HCPEB. Due to the low solubility of Zn and Cu and the extremely heating and cooling process of HCPEB, the Pb vapor deposited and crystallized on the top surface at nanoscales and then the special surface morphology was formed (Figure 5(b)). Thus, vaporization behavior induced by HCPEB treatment could lead the metal to produce fine grains at nanoscales with low boiling point.

Figure 5:

SEM images of nanocrystalline on the HPb59-1 alloy surface after HCPEB treatment with 15 pulses, 20 kV (a) Pb at low resolution; (b) Pb at high resolution; (c) two phases at low resolution and (d) two phases at high resolution.

Based on Cu–Zn phase diagram, the white substance was α-phase, while the gray matrix was β-phase, as shown in Figure 5(c) and (d). They are uniformly distributed with a grain size between 20 and 50 nm. The generation mechanism of this morphology could be attributed to the superfast melting and solidification process of HCPEB. The β-phase separated out firstly from the liquid and then α-phase was formed in β-phase with the decreasing of temperature. The α-phase could be retained which had no time to grow up on account of the ultrafast melting and solidification on the surface of the alloys, resulting in the formation of nanoscale structure of α-phase. The α-phase was uniformly distributed across the β-phase, which could play an indirect role in refining β-phase. Hence, a melted layer containing abundant nanostructured α- and β-phases was obtained by HCPEB treatment.

Figure 6 shows the surface grain-size statistics of Pb and α phase after 15-pulse HCPEB treatment. The average grain sizes of Pb and α were about 67.17 and 36.69 nm. They reached nanoscales. These nanocrystalline structures could also be observed after HCPEB treatment with 5 and 25 pulses, which was also reported in many references [14, 15]. Nanomaterials are typically characterized by ultrafine grains, which have special physical, chemical and mechanical properties for their superfine size and high density of grain boundaries. For practical engineering materials, it is an effective way to enhance the performances of wear resistance, corrosion and strength and so on, and thereby extending the service life of materials if nanocrystalline structures can be introduced in the serving surface of the workpieces. Consequently, HCPEB technology is a promising method for introducing nanocrystals in the near-surface layers of metallic materials.

Figure 6:

Surface grain-size statistics of α phase and Pb after 15-pulse HCPEB treatment.

Phase transition on the surface of alloys

XRD analysis showed that the initial structure of lead brass HPb59-1 was made of α, β′ and Pb phases (Figure 7). Compared with the original sample, a new phase (β) was found after HCPEB irradiation. When the alloy was cooled to 456–468℃ from high temperature, the solid solution of disordered β-phase would transform into ordered β′-phase. Presumably, the disordered β-phase could be preserved which had no time to convert into ordered β′-phase because of the ultrafast melting and solidification process. Therefore, we determined that the new phase was disordered β-phase by phase analysis of XRD.

Figure 7:

XRD diffraction patterns of HPb59-1 alloy surfaces before and after HCPEB treatment.

Phase transition can be divided into first- and second-order phase transitions thermodynamically. The transition between the ordered and disordered states belongs to the second-order phase transition, in which process there is no change of entropy and volume, but heat capacity and the coefficient of compressibility and expansion were changed. The phases of β and β′ have the same crystal structure of basic knowledge in the field of Crystallography (bbc). This transition is done by rearranging the atoms through the way of diffusion. The disordered β-phase possesses better plasticity compared to the ordered β′-phase, which is conducive to increase plastic property of alloy and expand the range of applications [16]. By contrast, we could also find that the diffraction peaks of Pb, α-phase and β-phase obviously broadened, which indicated that the size of grains was refined corresponding to the nanostructures observed through SEM after HCPEB treatment. In addition, the intensities of the three phases decreased significantly. It suggested that different elements in the alloy were dissolved more completely, which resulted in the formation of metastable structure – supersaturated solid solution.

Changes of residual stress and grain orientation

Compared with the initial sample, the diffraction peaks shifted toward high angles after HCPEB treatment, which illustrated there was a large residual stress on the alloy surface. The macroscopic residual stress in materials can be estimated according to the XRD peaks [14]:

where E and ν represent the modulus of elasticity and the ratio of Poisson, respectively. As for lead brass, E = 105 GPa, ν = 0.324. In the alloy of lead brass, β′-phase with poor plasticity has the tendency of stress concentration, so we take β′-phase as an example to calculate the residual stress. Assuming that there was the (110) direction of β′-phase grains perpendicular to the direction of thermal stress, grains remained in the stress-free state with spacing of d₀ before treated. Under effect of stress σ (110), the original d₀ turned into d_n after HCPEB treatment. According to the formula, we could figure out the residual stresses on the alloy surface with different pulses: residual stress σ= –1.023 GPa with 5 pulses, –1.864 GPa with 15 pulses and –1.425 GPa with 25 pulses. It was not difficult to explain why the microcracks appear on the surfaces of samples, which was due to the high residual stress beyond the ultimate strength of materials. Meanwhile, residual stress decreased with the rising of HCPEB pulses. This was due to releasing of part residual stress in the form of cracks when the surface residual stress accumulated to a certain degree. HCPEB treatment also caused the change of peak intensities of each crystal plane in the α- and β′-phases. These changes in diffraction patterns before and after HCPEB treatment could be explained by the superfast cooling and recrystallization induced by HCPEB process.