Laser Cladding of Mg20Al80 Powder on ZM5 Magnesium Alloy

By Chen, C J Wang, M C; Wang, D S; Jin, R; Liu, Y M

Laser surface engineering can be used to alter the surface of alloys and produce a fine microstructure, usually improving the corrosion resistance and wear resistance. In the present paper, a CO2 laser was used to melt preplaced Mg20Al80 powder with the objective of improving the corrosion resistance of ZM5 magnesium alloy in sea water solution. The corrosion behaviour of the ZM5 substrate and the laser cladding Mg20Al80 sample was studied with relating to their microstructures. Corrosion was studied in a 3.5 wt- %NaCl solution at pH 7 by observing the corrosion morphology. The results showed that the Mg^sub 17^Al^sub 12^ acted as galvanic cathode in both the ZM5 substrate and the cladding layer. The enhanced corrosion resistance was attributed to the fine dispersion of intermetallic phases and extended Al solid solutions by laser rapid solidification. Keywords: Laser cladding, Magnesium alloy, Galvanic corrosion

Introduction

Magnesium alloys have become candidate materials for many applications in industries. However, magnesium is vulnerable to galvanic attack. The improvements in corrosion resistance imparted to magnesium alloys by rapid solidification (RS) processes have been studied by many researchers.1-3 From an investigation of binary RS magnesium alloy ribbons, Makar et al. ‘ found that aluminium is the only element to cause a significant decrease in the corrosion rate of magnesium alloys. They showed that >10%A1 in the solid solution is favourable to the formation of passive films and that rapid solidification ameliorates the repassivation behaviour of Mg-Al alloys, this repassivation being superior for the alloys containing more aluminium.1″3

Although laser cladding of substrates alleviates corrosion losses with Al containing cladding offering the most improvement, the interface between the cladding zone and the substrate remains prone to accelerated corrosion. Furthermore, Al powder is, like Mg, highly active and pyrophoric; however, any reaction between the melted Al and the atmosphere can be avoided to a large extent by the rapid solidification of the Mg-Al melt. At the same time, the shielding gas argon should prevent the Mg-Al melt from being oxidised efficiently. Laser surface melting/cladding/ alloying are advantageous processes to improve the corrosion resistance of magnesium alloys.4″* Laser cladding produces surface layers with a fine microstructure that reduces the size of galvanic couples and expands the solid solution range of alloying elements. These factors can potentially improve the overall corrosion resistance.6″9 The present paper reports laser cladding of Mg-Al powders, and the characteristic microstructure and corrosion properties of the clad alloy.

Experimental

Three groups of binary alloys (Mg20A180, Mg50A150 and Mg80A120) were chosen for laser cladding experiments. The powder size and purity were:

The powder for laser cladding was placed on the ZM 5 surface using cellulose acetate with a thickness of ~2 mm. Laser cladding was carried out using a 5 kW continuous wave CO2 laser with a beam diameter of 3 mm and set at an output power of 1.5 kW. Argon gas was used as the shrouding environment (at a flowrate of 10L min^sup – 1^, which was delivered to the melted pool by a 4 mm diameter copper pipe) to avoid oxidation during laser cladding. The surface scan was ~ 10 mm S^sup -1^ and the overlapped ratio was about 20-40%. The focus length was ~ 300 mm. Figure 1 shows the schematic diagram of the experimental set-up used for laser cladding in the present study.

Generally, the Mg20A180 cladding sample had less porosity than other alloys. For magnesium rich (Mg80A120) laser cladding samples, the cladments are still thin because of the low molten cladment surface tension and high evaporation losses. Therefore, the observations reported here mainly focus on the Mg20A180 cladding sample.

1 Schematic diagram of experimental set-up for laser cladding of Mg-Al powders on ZM5

The microhardness of the coating and ZM5 substrate was measured using an HX-I Vickers microhardness tester with a 100 g load applied for 10 s. The microstructure of the coating was examined by optical and scanning electron microscopy (SEM). Polarisation curves of the treated and untreated ZM5 samples were measured in a 3-5 wt-%NaCl solution saturated with Mg(OH)2 using a PAR2263 electrochemical measurement system. A scan rate of 1 -66 mV s^sup -1^ was used in the measurements of the potentiodynamic polarisation curves. The reference electrode used in the measurements was a silver/silver chloride electrode and the counterelectrode was a platinum mesh.

Before the corrosion process, the samples were mechanically polished using SiC paper and then diamond paste, to obtain surfaces with the same morphology and the same roughness. Subsequently, the samples were cleaned for 3 min in an ultrasonic bath in pure acetone and dried. To determine the initiation process for the corrosion, a detailed analysis of the microstructure was undertaken. Metallographical samples were immersed in 3-5 wt-%NaCl solution and examined by SEM after the samples were taken out and dried. The total immersion time was 3 days and the pH was set to 7. After the corrosion procedure, the samples were rinsed in distilled water, dried and analysed.

For measurement of mass losses, samples were cut into approximately square pieces, cleaned in distilled water and dried immediately. After this procedure, the weight and superficial dimension of the samples were measured. The samples were then subjected to a 3-5 wt%NaCl salt solution test over seven days, with duplicate sample sets removed at intervals of 12, 24, 48, 72, 96 and 168 h. The corrosion products were removed in a chromic acid bath for 5 min. The samples were then rinsed in distilled water, dried and immediately weighed.

Results and discussion

Microstructure of coating and XRD results

Figure 2a and b show that the as received ZM 5 substrate has a two phase microstructure typically consisting of primary a grains surrounded by a eutectic mixture of a and beta phases (the intermetallic MgnAIi2). The lightest areas are beta-Mg17^sub ^Al^sub 12^ and the darkest areas are primary alpha-Mg. The phases between the beta-Mg^sub 17^Al^sub 12^ and the alpha-Mg are eutectic alpha- Mg, which is clearly shown in Fig. 2b. The aluminium is partly in the solid solution in the matrix and also precipitates in the form of Mg^sub 17^Al^sub 12^ along the grain boundaries as continuous primary phase as well as a secondary lamellar structure (see Fig. 2b). The compositions of alpha and beta phases are listed in Table 1 . Figures 2c, 3a and b show the cross-section of the laser cladding samples. Optical and SEM investigation of the laser clad samples showed that a complete fusion of the coating and the substrate material was obtained with no cracking visible at the interface. The thickness of the cladding layer was about 650-750 [mu]m. Thus, an adherent and defect free interface with negligible heat affected zone can be noted. Figure 3a and b shows two different microstructures obtained in the present study under the same laser processing parameters. A significant change in the solidification morphology occurs within the laser cladding zone by transition from cellular to dendritic, which is very evident from Fig. 3c and d.

2 Scanning electron microscopy Images showing ss microstructure of ZM5 substrate, b great detail of ZM5 substrate and c cross- section between coating layer and substrate

3 Scanning electron microscopy Images showing a, b cross- sections of coating layer and substrate, c, d microstructures of cladding layer with different morphologies and e, f detail microstructures of same field as c and d shown at higher magnification

Table 1 Summary of element analysis in regions marked as 1, 2, 3 and 4 shown in Fig. 3 and phases in ZM5 substrate*

The microstructure of the coating is found to vary with laser parameters and laser materials coupling,410 and mainly depends on the following three factors. First, solidification controls the size, shape and distribution of beta-Mg^sub 17^Al^sub 12^ phase in the final microstructure in the coating.” Second, the beta-Mg^sub 17^Al^sub 12^ exhibits a wide range of morphologies in hypereutectic Mg-Al alloys (in the present study, Mg20A180 constitutes a hypereutectic) depending on the composition and cooling rate.11 Third, the microstructure was formed by the rapid variation of GIR, where G is the thermal gradient and R is the solidification rate; the rapid variation of GIR is due to the instability of the laser energy. Therefore, the solidification rate and partition coefficient on the instability of the interface leads to the formation of different types of microstructures.12 The eutectic tends to become more cellular with increasing cooling rate, but more dendrite with lower cooling rate.11 4 X-ray diffraction patterns of a ZM5 substrate and b laser Mg20AI80 coating

Figure 3c-/ displays the microstructure of the cladding layer revealed by optical microscopy (OM) and SEM. The microstructure of the cladding layer is characterised by finer and denser structures than that of the ZM5 substrate. The precipitates are refined along the boundaries of the coating. At this magnification, the cladding zone was homogeneous compared with the ZM5 substrate. The microstructure of the laser cladding zone consists of fine dendritic grains (Fig. 3c) and cellular grains (Fig. 3d) growing epitaxially from the liquid/solid interface. It can be clearly seen that the substrate grains are largely (~ 10 times) coarser than in the laser cladding zone. A detailed composition analysis was undertaken by SEM- EDS with the grains and the precipitates for the coating layer. The element analysis of the selected regions marked as 1 , 2, 3 and 4 in Fig. 3e and f are summarised in Table 1 . It can be seen from Table 1 that for Mg20A180 powder, alpha-Al and beta-Mg^sub 17^Al^sub 12^ phases should be formed as predicted by the equilibrium phase diagram. However, alpha-Mg and beta-Mg^sub 17^Al^sub 12^ phases were formed according to the EDS results. The results imply that the total coating layer was really an alloying layer.

5 Microhardness profile along depth of Mg20AI80 coating and ZM5 alloy

Figure 4 shows the XRD patterns of the laser cladding Mg20A180 powder on the ZM5 substrate and the untreated ZM5 substrate. XRD analysis of the untreated ZM 5 substrate confirmed the presence of alpha-Mg and beta-Mg^sub 17^Al^sub 12^ phases. After laser cladding of the Mg20A180 powder, a large amount of alpha-Mg and beta-Mg^sub 17^Al^sub 12^ phases were produced in the clad layer. Aluminium particles in the powder were dissolved in the melt pools during laser cladding, which increased the content of Al in the liquid alloy and led to the formation of a solid solution of Al in the Mg phases during the rapid solidification. According to the XRD spectra shown, they are labelled the same two phases, but the peaks are not always in the same position. This situation is also caused by the extended Al solid solidification in the alpha-Mg and beta-Mg^sub 17^Al^sub 12^ phases, which changes the lattice parameters of the phases.

Microhardness

Figure 5 presents the microhardness profiles of the laser cladding ZM5 magnesium alloy with preplaced Mg20A180 powder. It is obvious that the microhardness of the coating has greatly increased to 320-350 HV as compared to the ZM5 substrate (96-100 HV). The higher microhardnesses are attributed to the grain refinement and extended solid solution due to laser rapid solidification.4″9

Corrosion tests

The variation of corrosion rate as a function of exposure time for the ZM5 substrate and the laser clad samples is given in Fig. 6. The ZM5 substrate showed an initial decrease in corrosion rate followed by an increase up to 4 days of exposure. However, in the case of the laser clad samples, the corrosion rate showed a steady decrease with exposure time except for 72 h where it showed an upward tread. The corrosion rate of the ZM5 substrate was slightly higher than that of the laser clad samples at all exposure times.

The polarisation behaviour of the laser clad and untreated ZM5 samples are shown in Fig. 7. The results indicated that the laser cladding Mg20A180 coating is superior to the untreated ZM 5 substrate in corrosion properties. The corrosion potentials and corrosion currents for the tested samples are obtained from the intersection of cathodic and anodic Tafel plots. The corrosion current of the Mg20A180 coating is of the same order of magnitude as that of the untreated ZM5 substrate. However, the corrosion potential of the laser treated sample was shifted to a nobler position (~ 170 mV) than that of the untreated ZM5 substrate, which is likely to increase the corrosion resistance of the treated sample.

6 Corrosion rates of a ZM5 substrate and b Mg20AI80 laser clad sample

7 Corrosion results of potentiodynamic anodic polarisation for Mg20AI80 coating and ZM5 alloy

8 a corroded surface of ZM5 specimen which shows selective dissolution occurred in surface, b and c corrosion morphologies of laser Mg20AI80 coating for 24 h immersion test

A detailed analysis of the microstructure of the corroded films was undertaken to investigate the initiation of corrosion. Figure 8 shows SEM images of the top surface of corroded samples of both ZM 5 substrate and laser clad samples, after they were immersed in the 3- 5 wt-%NaCl solution for 24 h. It is apparent from Fig. 8a that corrosion starts in the regions along the beta phase edges, rather than wholly within the beta phase or the a matrix. It is relevant to mention that the microstructure of the ZM5 specimen mainly consists of a grains with the beta phase (the intermetallic Mg^sub 17^Al^sub 12^) along the a grain boundaries (Fig. 2a and b). In this regard, it is relevant to mention that the beta phase, which shows considerably more corrosion resistance, is cathodic to the alpha phase and hence, the two phases in contact are likely to result in microgalvanic corrosion.13″17 As a result, corrosive attack starts at the phase beside these beta phase edges as shown in Fig. 8a.13- 17

9 Scanning electron microscopy Images showing a,b corroded surface of untreated ZM5 sample, c, d corroded surface of laser treated samples, e, f detailed microstructure of corroded surface of laser treated samples as shown In c and d for 72 h immersion test: A, C indicate zones corroded and B, D indicate zones not attacked

On the other hand, the laser clad samples remain mostly unaffected after immersion for the same time of 24 h (see Fig. 86 and c). By comparing Fig. 86 and c with Fig. 3c and d, it can be seen that corrosion was preferentially started from the white phases. Thus in Fig. 3c and d, the white zone can be clearly seen whereas in Fig. 86 and c, it is less clearly visible. Laser treatment can create a fine microstructure along with Mg17AIj2 intermetallic precipitates along the grain boundaries. The enhanced corrosion resistance of the treated samples can be attributed to the decreased anode to cathode ratio due to grain refinement and precipitation of Mg17Al12 intermetallics along the grain boundary.13- 17

Figure 9a and b shows the SEM images of the corroded film on the untreated ZM5. Figure 9a shows both the alpha grain interiors and its grain boundaries whereas Fig. 96 shows only the a phase beside the beta phase edges corroded, i.e. the same situation as in Fig. 7a, for 72 h immersion. In ZM5 alloy, the alpha grains should be highly polarised because the beta phase acts an effective cathodic component along their boundaries. Thus, in the test solution, the eutectic chi at the grain boundaries should be corroded first, as indeed shown in Fig. 8a. As to Fig. 9a and b, it is easy to imagine that the corrosion was initiated by the dissolution of the eutectic alpha along the grain boundaries in the vicinity of the beta phase. After most of the eutectic alpha along the grain boundaries were removed by dissolution, some small and discontinuous beta phase particles, which were surrounded by the eutectic alpha, might fall out by undermining. Subsequently, the primary a of the grain interior was corroded.13-17

However, the film is continuous in the laser clad samples with the omission of the white phase on the surface. The presence of precipitate phases could be detrimental to the corrosion resistance of magnesium based alloys.16 In the present study, the fine beta- Mg17Al12 intermetallic phase precipitated along the grain boundaries acting as cathodic to the grain and corrosion began from the beta matrix interface. Hence, the grain boundary and grain together act as a galvanic couple.916 From Fig. 9c-/, it can be seen that some of the white phase was missing after 72 h immersion. By comparing Fig. 3e and / with Fig. 9e and /, one can see that the white phases were almost completely missing after 72 h immersion (white phases are indicated by 2, 4 and A, C).

The predominant factors determining the rate of galvanic corrosion are the anode/cathode area ratio and the difference in potentials of the two phases.1516 The smaller the grain size, the higher the grain boundary areas. Therefore, more beta-Mg]7Al)2 precipitates and thereby, decreases the anode to cathode area. As a result, the corrosion rate is lower in laser clad samples. Furthermore, the extended solid solubility of Al in the Mg matrix, caused by the rapid solidification after laser cladding, greatly improves the overall general corrosion resistance of the matrix.9

Concluding remarks

Mg20A180 powder was preplaced on ZM5 magnesium alloys and remelted by the laser cladding technique. Al solid solution alpha- Mg and beta phases were obtained in the coating. The microhardness of the coating was greatly increased compared with the untreated ZM5 substrate. After laser cladding, the corrosion resistance was also superior to that of the untreated ZM5 substrate due to the refined microstructure, the decrease in the anode to cathode area and the extended solid solubility of Al in the Mg matrix. The primary corrosion mechanism is galvanic corrosion due to the Mg)7Al]2 acting as cathodic phase to the a phase.

Acknowledgements

The authors acknowledge the support by National Nature Science Foundation of China (50371093). The authors would like to express their appreciation to Dr Liu Jianguo, Zhang Min and Han Dongyun for their kind help in experiment tests.

(c) 2007 Instituto of Material- , Minorali and Mining

Published by Manoy on behalf of the Institute

Received 10 March 2006; accepted 15 August 2006

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C. J. Chen*1,2, M. C. Wang1,2, D. S. Wang1, R. Jin1 and Y. M. Liu3

1State Key Laboratory for Corrosion and Protection, Chinese Academy of Sciences, Shenyang 110016, China

2Graduate School of Chinese Academy of Sciences, Beijing 100039, China

3Shenyang Liming Aero-engine Corporation, Shenyang 110043, China

* Corresponding author, email chjchen2001@yahoo.com.cn

Copyright Institute of Materials Jun 2007

(c) 2007 British Corrosion Journal. Provided by ProQuest Information and Learning. All rights Reserved.