Surface Engineering of Magnesium Alloys: Surface Modification to Improve Corrosion Resistance

Magnesium (Mg) and its alloys have been the topic of intense research over the past 15 years as the automotive and aeronautic industries strive to increase fuel efficiency by reducing the weight of vehicles. Widespread implementation is currently limited by the poor intrinsic corrosion resistance of Mg and its alloys stemming from the lack of a passivating film and their high susceptibility to galvanic corrosion. Poor intrinsic corrosion of Mg alloys is often attributed to heterogeneous microstructures containing electrochemically noble secondary phases which promote the preferential dissolution of the Mg matrix. The research herein is aimed at the mitigation of localized corrosion which stems primarily from the secondary phases present in Mg alloys by non-equilibrium surface processing.
The initial thrust of the research was intended to describe controlling factors for localized corrosion (micro-galvanic couples) which occur during full immersion of specific Mg alloys (AZ31B-H24, AZ91D, and AM60B) in NaCl solutions. A focus was placed on the critical size of an electrochemically noble secondary phase particle required to induce localized corrosion, breakdown of the partially passive film, or the initiation of filiform-like corrosion (FFC). When immersed in a stagnant 0.6 M NaCl solution, the AZ31B alloy exhibited local pitting-like corrosion for the areas surrounding β-Mg17Al12 (radius < 250 nm) particles while the initiation of FFC was most often observed in proximity to the larger Al-Mn secondary phase particles (radius > 250 nm). The smaller Al-Mn secondary phase particles (radius < 50 nm) showed minimal contribution to the local corrosion processes.
The second thrust investigated the near surface dissolution of the Al-Mn and Mg-Al secondary phase particles present in the Mg alloys, utilizing a non-equilibrium surface process with a pulsed ultraviolet excimer laser. Analytical calculations showed that after a single laser pulse, and a typical melt time of ~50 ns, secondary phase particles with a radius of 75 nm could be dissolved. However, experiments utilizing scanning electron microscopy (SEM) observed complete dissolution for particles up to a radius of 250 nm, within the resolution of the SEM (~5 nm), and partial dissolution for the larger Al-Mn particles. The additional homogenization afforded by excimer laser processing was explained by the liquid mixing generated by a plasma pressure wave and the low melting temperature of the β-Mg17Al12 particles (Tm = 723 K) which led to complete melting. The laser induced plasma formed above the substrate during processing was found to generate a pressure of ~50 MPa which acted on the melted substrate surface, resulting in material transport of several µm, experimentally shown to be ~3 µm per pulse. This material transport could also push the liquid Mg on top of the larger Al-Mn particles, minimizing the particles impact as local cathodes during corrosion. Overall, when a single laser pulse was applied to AZ31B, the density of secondary phase particles was reduced by approximately an order of magnitude. This reduction in secondary phase particle density led to a reduction in cathodic the half-cell hydrogen evolution reaction (HER) kinetics for all Mg alloys. The laser processing parameters were
optimized around this corrosion response ultimately yielding an order of magnitude reduction in cathodic HER kinetics (fluence = 1.5 J/cm2, pulse per area = 200, pulse overlap = 50%, and Ar pressure = 810 Torr).

The third thrust investigated the full immersion corrosion response for the laser processed Mg alloys in a stagnant 0.6 M NaCl solution. It was determined that the order of magnitude reduction in cathodic kinetics resulted in an order of magnitude increase in polarization resistance (inversely proportional to corrosion rate) for all Mg alloys. The polarization resistance for AZ31B-H24, AM60B, and AZ91D began at 200 Ω-cm2, 2,000 Ω-cm2, and 2,000 Ω-cm2, respectively and after laser processing were increased to 5,000 Ω-cm2, 23,000 Ω-cm2, and 28,000 Ω-cm2, respectively. The open circuit potential (OCP) was more negative for the laser processed Mg alloys, similar to high purity Mg, and observed a gradual increase to values similar to the as-received alloy. The initiation of FFC occurred shortly after the OCP of the laser processed specimen reached the OCP of the as-received specimen, typically after 24 to 60 hours of full immersion. This extension in the time to initiation of FFC is substantial considering the relatively aggressive corrosion environment (full immersion in 0.6 M NaCl), and should observe an additional increase in time to initiation in more applicable corrosion environments, such as field exposures in non-marine environments.

The use of x-ray diffraction (XRD) and Fourier transform infrared (FTIR) spectroscopy determined the corrosion product for all specimen to primarily consist of Mg(OH)2 and a layered double hydroxide (LDH) phase (likely hydrotalcite). LDHs have been observed on Al containing Mg alloys previously;
however, the ability for a LDH corrosion product to act as a passive layer is often short lived because the initiation of FFC occurs quickly in aqueous NaCl solutions, dominating the electrochemical response of the alloy. Interestingly, the LDH formed to some extent on all specimens after immersion in the stagnant 0.6 M NaCl solution, regardless of the heterogeneity of the Al containing Mg alloys microstructure. The difference in corrosion product for an as-received and laser processed surface was based on the quantity of Mg(OH)2, with the as-received alloys exhibiting more Mg(OH)2 in comparison to the laser processed specimen, indicative of the higher coverage of FFC on the as-received alloys. That said, the observation of the LDH phase on all specimen may have a wide spread impact on the corrosion analysis of Al
containing Mg alloys since the primary corrosion product at short immersion times for these alloys was verified to be the LDH phase which precipitates at a pH of ~8 in contrast to Mg(OH)2 precipitating at a pH of ~10.5.

The fourth thrust of this thesis was centered on the deposition of a multilayered coating composed of MgO/Mg(OH)2 and Gd(OH)3 by pulsed laser deposition onto the AZ31B-H24 surface. The 300 nm thick films were successfully applied and acted as a passive coating, inhibiting the initiation of FFC on the as-received and laser processed AZ31B-H24 alloy for 1 and 72 hours, respectively. Increasing the initiation time of FFC to 72 hours in 0.6 M NaCl is similar to results for hydrotalcite conversion coatings on Mg alloys. The polarization resistance of the laser processed and coated specimen increased to a value of ~100,000 Ω-cm2, decreasing with immersion time to ~30,000 Ω-cm2 until the initiation of FFC occurred after 72 hours. Initiation of FFC on the laser processed and coated specimens was preceded by a “repassivation” like phenomenon where local defects in the coating were filled in by the conversion/precipitation of Mg(OH)2.

The fifth and final thrust of this thesis investigated the specimen’s corrosion response to a B117 salt spray exposure as a harsh, accelerated outdoor exposure. A LDH phase was formed on all specimens, similar to the full immersion conditions, with the only detectable corrosion product found on the laser processed specimens being the LDH, signifying the lessened coverage of FFC for the laser processed surfaces. In addition, the LDH phases were found to be in competition for the most thermodynamically stable corrosion product with the carbonate rich phase, hydromagnesite, with experimental results verifying the presence of Al suppressed its formation. Regardless of the corrosion products formed, the laser processed specimens consistently observed slower corrosion rates than their as-received counterparts, albeit with a less dramatic reduction in corrosion rate, only ~3 fold for the salt spray versus ~10 fold for the full immersion experiments.