11 L GRANASY Trans. Jap . Inst. Metals , 1986, 27, 51-60. 12: L:D. LANDAU and E.M. LIFSHITZ, Statistical Physics , Pergamon , oxf~r~~;~5~980 . 13 S PATANKAR Numerical Heat Transfer and FLuId Flow , HemIsphere, Ne h' 1987 14: S:N.ZOLOTAREVandA.N . SHUMAKOV, Physics of Metals and Metallog rap y, ,
15 f't/~I~~ENT J .G . HERBERTSON and H.A. DAVIES, in Rapidly Quenched Metals . TV' • d T Masu'motoandK . Suzuki , The Japan Inst. Metals,Sendal , 1982, Vol. 1, pp. 77-,e s. .
80.
(Received 18th January 1990 and accepted 12th May 1990)
The Sensitivity to Iron Impurity Content of the Corrosion Rate of Extrusions of Rapidly-Solidified Mg-15wt% AI Alloy Powder
J.D. Cotton* and H. Jones
School of Materials, University of Sheffield, Mappin St., Sheffield S1 3JD
Abstract The free corrosion in 3wt% NaCI aqueous solution of extrusions produced from rapidlysolidified (RS) atomized Mg-15wt%AI alloy powder has been studied by hydrogen evolution and weight loss measurements. Trace levels of Fe in the alloys , varying from 0.003 to 0.020wt% , were demonstrated to have a strong exponential effect on the instantaneous corrosion rate. However, the sensitivity to Fe content for the RS alloys in this study is shown to be markedly below that of previous studies for conventionally cast Mg-AI alloys, which utilized much lower solidification rates . This effect of rapid solidification rate is explained on the basis of the decrease in scale of microstructure it produces. Both pitting and filiform modes of corrosion were identified and characterized by optical and scanning electron microscopy. Brucite was identified as the dominant corrosion product in addition to another compound which may be a modified form of hydrotalcite .
1. Introduction Magnesium and its alloys have long offered great potential for structural weight reduction in the aerospace and automobile industries, but have been held back by poor corrosion resistance and low formability [1]. Due to its extreme position in the galvanic series, in conjunction with an easily perturbed surface film, the corrosion resistance of Mg is strongly linked to galvanic microcells set up between microstructural constituents of more noble potential. In particular, heavy metal impurities such as Fe, Ni and eu are highly deleterious. Thus, two avenues have been taken to improve the corrosion resistance of Mg: the reduction of heavy metal impurity levels [1] and microstructural refinement [2] . Refinement may be accomplished by
* Present address: Department of Materials Science and Engineering, University of Florida, Gainesville, Florida , USA.
56
rapid solidification processing (RSP), which improves the corrosion resistance by reducing the scale of the microstructure and effectively distributing the corrosion more evenly. In addition, RSP may increase the extent of solid solubility which can be important for the creation of stable, passive films or allow the use of more active (relative to Mg) alloying elements, such as certain rare earths.
Active metals, such as magnesium and its alloys, are subject to pitting corrosion due to the presence of discontinuities in the protective surface film. Sources of film discontinuities are many, varying from emergent screw dislocations to macroscopic inclusions [3], some requiring time to activate before pits are observed [4]. Impurity phases of metals such as Fe, Ni and Cu are ideal pit initiation sites because of the associated break in the protective film combined with the low hydrogen overvoltage of the impurity phase . Microgalvanic corrosion between Mg and these metals is spontaneous in aqueous solution and accelerated by the presence of chloride ions [5]. Once initiated, pit growth is auto-catalytic and occurs at an accelerating rate, with the locally acidic environment preventing repassivation within the pit. Hydrogen gas is produced at the cathodic sites and can be collected as a method of determining corrosion rate by stoichiometric calculation with the overall corrosion reaction [4,5]:
Mg + 2H20 = Mg(OH)z + H2 (1)
Note that, within the pit, the presence of a high concentration of chloride ions prevents proton recombination initially and produces a locally acidic solution.
Hydrogen evolution has been used previously [6-9] as a technique for monitoring the corrosion of active metals, particularly Mg. One aspect of this method that is quite intriguing is the observed 'negative difference effect' in which the corrosion rate measured by hydrogen evolution is less than that measured by weight loss. This effect has been ascribed to a number of causes, including the loss of metallic Mg particles, monovalent Mg, undermining and loss of heavy metal impurity particles, and film damage and repair processes [6,10-15]. Since monovalent Mg has not been observed, undermining of impurity particles would not necessarily produce a large difference in weight, metallic Mg particles would not be stable (would eventually corrode), and film repair is not important in stable pits, these explanations are less than satisfactory. Two alternative explanations are offered within the context of the current study. One source of the difference is the weight loss contribution from second phases, such as Mg17AI12' which are stable when exposed to the solution. These phases are ultimately lost with the corrosion product. The second is that, during pitting
57
corrosion, protons are produced within the pit , many of which combine to form hydrogen gas. However , as the pit grows a significant proportion will remain as H + in response to the high CI - concentration within the pit , and are responsible for the increased acidity . Thus, the proportion of protons which do not eventually combine to form hydrogen gas will cause an apparently lower corrosion rate when measured by hydrogen evolution.
Early studies of the corrosion resistance of Mg and its alloys in aqueous environments produced very erratic results . In a classic study, Hanawalt , Nelson and Peloubet [16] found the inconsistencies to be a result of extremely small amounts of impurities , one of the most insidious being Fe . They determined a 'tolerance limit' for Fe in unalloyed Mg of 0.017wt%, above which the corrosion rate increased markedly beyond the passive rate of 12 mils per year. The addition of lOwt% AI to the Mg decreased this limit to 0.0002wt% Fe or lower. In addition, Mn was shown to have a beneficial effect on the Fe tolerance limit, at concentrations of 0.2wt%Mn for high AI alloys. No Mn additions were made to the alloys in the current study. More recent work by Hillis [1] on die-cast material suggested that the sensitivity to Fe is not as large as the Hanawalt study indicated. This discrepancy may be due to the difference in coarseness of microstructure between the mold-cast alloys in Hanawalt 's study and the die-cast alloys in Hillis' study. Given this effect of microstructural scale, it is expected that further refinement , such as that derived from RSP , will increase the tolerance limit for Fe still more. This has been observed in other work on Mg alloys [6] and appears to be the case in the current study in which the tolerance limit is not well defined and is well above Hanawalt's limit of 0.0002wt%.
The results presented here compare the sensitivity to Fe content of the corrosion of RSP Mg-AI extrusions to that of conventionally-cast Mg-AI , alloys. The extrusions were produced from atomized powder containing differing Fe levels , between 0.003 and 0.020wt%. The lower sensitivity to Fe content for the RSP product is explained in terms of microstructural refinement. In addition, the corrosion forms , rates and products are characterized.
2. Experimental Procedure The Mg-15wt%AI extrusions used in this study were produced in a prior research program [17] by extrusion of RSP atomized powder. AZ31HP, an experimental high purity version of AZ31B, was used as comparison material known to display good corrosion resistance in wrought form. The composition of each of the extrusions, including analyses for Fe and Ni, are shown in Table 1. The Mg-15wt%AI extrusions were provided as rod 10 mm in diameter and the AZ31HP as extruded plate 10 mm in thickness . The amount of oxygen present in the atomizing gas during powder production
58
TABLE 1
Results (wt%) of wet chemical analyses of alloys
Element Extrusion-A Extrusion-B Extrusion-C AZ31HP
AI 14.5 14.5 15 .0 2.9 Zn nfa nfa nfa 0.69 Si nla nfa nfa 0.01 Mn nla nfa nfa 0.23 Fe 0.003 0.009 0.020 0.004 Ni 0.001 0.002 0.002 0.005 Sn nfa nfa nla 0.05 Pb nfa nfa nfa 0.03 Cd nfa nfa nla 0.03 Mg bal bal bal bal
Data provided by T . Wilks, Magnesium Elektron Ltd, UK. nla = not analysed. . Added oxygen contents of the atomizing gas used to produ~e powders for extrusIOns A , Band C were 1.0 percent, nil and nil respectively with 1.0, 011 a.nd 0.1 percent added oxygen , respectively in the collector system. Classification of resultmg powders was done m argon without added oxygen.
was varied as shown in Table 1. Such low levels of oxygen were necessary to allow safe handling of the powder.
The corrosion rate of each extrusion was determined by measuring the rate of hydrogen gas evolution from a specimen when immersed in 3wt%NaCI aqueous solution saturated with Mg(OH).z' All tests . were conducted at room temperature (22 ± 3°C). The corrosIOn test specImens were cut into shapes with either square, circular or semicircular cross section, depending on the product form. All surfaces were then ground on 600 grit silicon carbide paper, thoroughly rinsed in met~anol , followed .by drying, weighing and dimensional measurement: HandlIng w~s done Wlt~ plastic tweezers and the specimens were stored In sealed plastIC bag~ untIl tested. During corrosion testing the specimens were suspe~ded In the solution by polypropylene clips fashioned to contact ~he specu~en. at two opposing points. The hydrogen evolved from the specImen dUrIng Immersion in the solution was collected in a calibrated burette. Measurements were made hourly until the corrosion rate stabilized. Selected specimens were later stripped on their corrosion product by immers~on in boiling 1~% chromic acid containing 1 % silver nitrate and then rinsed In tap water , drIed and weighed to determine the corrosion rate by weight loss. .
By filtering the used NaCI solution following a tes.t, t~e Ins.oluble corrosion product of each extrusion was collected .. FollOWIng aIr ~ryIng for several days , this product was analyzed by X-ray dl~~actometry wlt.h Co Kalpha radiation using standard techniques with a PhIlIps PW 1700 dIffractometer. The ASTM procedure for the identification of compounds was used.
59
The microstructure of each extrusion was studied in the un etched condition. Standard metallographic techniques were used with the exception of the use of MgO powder in deionized water for the final polishing step. Selected corroded specimens were mounted in transparent epoxy, ground and polished. Photomicrographs were taken at several different magnifications of the metal/corrosion product interface. By placing small drops of the NaCI solution on a polished specimen surface and focussing through the solution, in situ observation of the corrosion process was carried out and recorded photographically.
A CAMSCAN Series 2 scanning electron microscope (SEM) , in conjunction with a LINK Systems X-ray Digital Analyzer, was enlisted to find and identify impurity phases that could have been responsible for the large differences in corrosion behaviour between the three Mg-15wt%AI extrusions. The method adopted to locate these phases was to allow a polished surface to corrode in the salt solution for approximately 30 minutes, rinse in methanol and dry. Cathodic particles were then located in the secondary electron imaging mode by a 'halo ' of corrosion debris. Photomicrographs were taken at several magnifications in both secondary and backscattered modes and X-ray energy spectra were recorded for selected areas of each microstructure to characterize the various phases present.
3. Results
3.1 Corrosion Rates
Plots of the corrosion rates (average of two or three specimens per condition) determined by hydrogen evolution are shown in Figure 1. The corrosion rates were evaluated, assuming an alloy density of 1.8 X 103 kg/ m3
, from the equation:
C.R. (mpy) = 2150.64(v/AT) (2)
where V is volume of hydrogen evolved in ml , A is surface area of specimen surface in square centimetres and T is immersion time interval in hours. Equation (2) is based on the overall corrosion reaction (1). Study ofthe plots reveals substantial differences in the steady state corrosion rates of the three Mg-15wt%AI extrusions. These are shown in Table 2. The results are probably affected by the formation of a proportionate amount of AIcontaining corrosion product. Due to its higher valence, aluminum can form 1.5 moles of hydrogen per mole of metal , and the apparent corrosion rate predicted by equation (2) may be slightly higher than that for unalloyed magnesium.
60
a a N
Q) U. oe. C') a a a
«
Q) a;-u. ~ u. 0
~ OJ 0 a .,. a a 0 a
2-CD c..
::c a ~
a «
en E. CII E t= c 0 'iii ... CII E .s
C .9 ; '0 '" 0 '" Z ~ 0 «)
"0
" .... ~ :>
.D
.= " .§ !: 0 .~
" E .§ '0 !: 0
"B !: :> ""' co '" '" !: .9 ; '0 > " !:
" 0/) 0 ....
"0 >-..c: >-
.D "0
" !: .§ .... " ., "0
'" '" '" B e !: 0 'v; 0 .... :5 u
'" :> 0
" !: co t: 19 '" ~
61
TABLE 2
Measured stable corrosion rates (mils per yr) of RSP Mg-15wt% AI and wrought ingot AZ31 HP alloys (mpy)
Extrusion Instantaneous (hydrogen)
EWLCR (hydrogen) True (weight loss)
Extrusion-A Extrusion-B Extrusion-C AZ31HP
4 178
2267 5
4 167
1691 10
7 196
2030 13
By measuring the cumulative volume of hydrogen evolved from each specimen, the equivalent weight loss corrosion rate (EWLCR) could be determined for proper comparison with actual weight loss data. Plots of the EWLCR are shown in Figure 2 with the points for the weight loss data. The actual (true) weight loss corrosion rates, determined gravimetrically, were evaluated from:
EWLCR (mpy) = 191601O.S(W/ST) (3)
where W is weight loss in grams, S is surface area in square centimetres and T = total immersion time in hours .
The closed point symbols in the plot are actual weight loss data for comparison .
The effect of Fe content on the corrosion rate is shown in Figure 3 by plotting the true weight loss corrosion rates of the three Mg-1Swt%AI extrusions against Fe content. The three data points indicate an exponential dependence on Fe content. Figure 3 also includes the data of Hanawalt [16] for mold-cast Mg-lOwt%AI alloys. The fitted curves show a difference in corrosion rate of at least two orders of magnitude between the two sets of results for Fe contents less than O.OOSwt%.
3.2 Identification of Corrosion Products: Mg(OH)z (brucite) was identified as the dominant corrosion product for all alloys tested and the only corrosion product for AZ31HP. X-ray peaks corresponding to MgJ7AIJ2 were also identified in the corrosion product of the RS Mg-1Swt%AI extrusions. Diffraction peaks from an unidentified third phase were also observed, corresponding to the following d-spacings (Angstoms): 8.1S, 4.09, 2.60 and 1.S3 . These displayed some degree of peak broadening. It is likely that other peaks of low intensity belonging to this phase were present but these coincided with those of the other compounds. Positive identification of this phase was not achieved, although the ASTM
FIGURE 2 Effect of immersion time on equivalent weight loss corrosion rates as determined by hydrogen evolution (solid lines) with symbols representing data from direct weight loss experiments.
FIGURE 3 Effect of iron content upon true weight loss corrosion rate for RS Mg-15wt. % Al (present work) compared with results of Hanawalt [4] for conventionally cast Mg-lOwt. %Al.
0\ tv
0\ (;.l
64
search manual showed it to closely match those of Mg6AI2C03(OH)16.4H20 (hydrotalcite) . The match was generally good , except that the c-dimension of the hexagonal hydrotalcite unit cell appeared slightly compressed (by about 5%) compared to the tabulated value. This may be a result of the substitution of chloride ions for hydroxyl ions in the lattice , which has been similarly suggested for Mg(OHh [5, 18].
Cathodic second phases in the microstructures were identified by X-ray diffractometry of saw filings from each alloy. The data revealed only metallic magnesium in AZ31HP and metallic magnesium plus Mg17A11 2 in the RS Mg-15wt% AI extrusions . Intensity calculations , based on the peak areas , showed that the relative volume fractions of Mg17A11 2 in the Mg-15wt% AI extrusions were within 10 percent of one another .
3.3 Light Optical Microstructures The longitudinal microstructures of the extrusions are shown in Figures 4ac. The scale and character of the RSP extrusions are similar in all but two respects. The first is the amount of interparticle oxide, which appears dark in the micrographs . Extrusion A contained the largest amount of oxide while extrusion B contained the least which corresponds to the amount of oxygen present in the argon atomizing gas. The oxide level appeared not to affect the corrosion rate . The second is the presence of hard , spherical particles in extrusion B (Figure 4b). Previous analysis showed these particles to be enriched in aluminum [17], and it is likely that they are powder particles containing a somewhat higher volume fraction of Mg17AI12 than the remaining matrix . Poor mixing in the crucible could have caused a small proportion of the atomized powder to have a higher AI-content , with only a small reduction in the AI-content of the bulk. The particles, in some cases , exhibit substantial plasticity, being elongated by the extrusion process (Figure 4b) . It has been shown [19] that unalloyed Mg17AI12 is very brittle below about 400°C. Since the extrusion took place at 200°C or less , the particles would have fractured if they were composed entirely of Mg17AI12 ' as was observed in the AZ31HP samples .
3.4 Metaliographic Sections of Corroded Specimens The sections failed to reveal any microstructural feature causing preferential corrosion for any of the extrusions . Nonetheless, there was a lack of attack of the spherical particles in extrusion B and they were observed to be spalled unaltered with the corrosion product in several cases (Figure 5) . This indicates that the AI-content in the particles may have been high enough to induce passivation by a compound more stable than the usual Mg(OHhRecent work [2, 20] has shown that AI-contents in excess of approximately 30wt% induce very good corrosion resistance . While pitting was the primary
65
(a)
(b)
(c)
FIGURE 4 Light optical microstructures: (a) Mg-15wt. % Al-O.003wt. %Fe (extrusion A), (b) Mg-15wt .% Al-O.009wt. % Fe (extrusion B) and (c) Mg-15wt. % Al-O.020wt. %Fe (extrusion C).
66
FIGURE 5 Hard spherical particle spalled with corrosion product without significant corrosive attack.
mode of corrosion in all specimens, those with the lowest Fe content, AZ31HP and extrusion A, experienced less severe attack.
3.5 In Situ Observation of Corrosion The corrosion of the extrusions was observed to proceed by two different modes, pitting and filiformt. These are shown in Figure 6. Pitting was indicated by slow hydrogen bubble growth and discharge at specific sites in the microstructure, although the exact nucleation sites were not identifiable. The'bubbles would nucleate very quickly (within 10 seconds), slowly grow to a critical size in the course of minutes, and detach, after which the process would repeat. At the most active sites, usually an Fe-rich particle, a 'halo ' [21] of corrosion product and debris was observed to develop around the site. The corrosion product was usually fractured (Figure 7), and was found to be enriched in Mg and chlorine by SEM analysis. Upon rinsing, the sites appeared light gray to the naked eye. The filiform corrosion was characterized by small, meandering trenches, in which the head would rapidly
t The filaments were not filiform corrosion in the classic sense in that no surface coating was present and it occurred in complete immersion. Filiform corrosion is usually defined as occurring in humid air beneath coatings (viz. paint) [22]. Its formation on bare metal exposed to bulk solution has scarcely been reported [3 , 10] and it seems likely that a different mechanism may be operative.
FIGURE 6 In situ observation of corrosion of extrusion B: (a) 15 seconds exposure to solution , (b) 210 seconds exposure. Dark circles are hydrogen bubbles which are slowly growing. Wavy features are "filiform" corrosion .
68
FIGURE 7 Scanning electron micrograph of corrosion product imbedded in a filament trail.
evolve tiny hydrogen bubbles in a continuous stream and leave behind a trail of imbedded corrosion product. Figure 8 portrays this process as it occurred in extrusion B. It is also notable that the filiform corrosion occurred preferentially in the extrusion direction of the longitudinal section, perhaps finding propagation more easy along the interparticle oxide/metal interface. This was verified by placing an intentional scratch nonparallel to the extrusion direction and observing filiform branches to extend from the scratch in the extrusion direction. Filiform corrosion appeared to nucleate at any significant surface discontinuity, particularly the specimen corners.
3.6 Scanning Electron Microscopy : Although general surveys of the slightly corroded specimens by SEM revealed no Fe-rich particles in extrusions A or AZ31HP, a number were found on extrusions Band C. Subsequent SEMIEDAX analysis of the particles in extrusion B showed them to contain primarily Fe (with trace Mn , V, Cr or Ni) . Figure 8 shows a typical Fe-rich particle in extrusion B with results of EDAX analysis . Analysis of cathodic particles in extrusion C indicated they were also Fe-rich , Figure 9. This might be expected from the chemical analyses of the extrusions. In some cases, the size of the Fe-rich
69
(a)
UF
c U ~ c aC U 'i F }. I .~ ~ fI v ,':, !
t-......!!'-LJ!1\11\I...-~~. ~-'-'--::::. J;~~, .... .. . . ~ .' ....... ..l "'\.. A 1+. 880 k eU 1 0 • 0 >
FIGURE 8 (a) EDAX spectrum of Fe-rich particle in extrusion B. (b) electron micrographs of partIcle (secondary electron Image on left, back-scattered image on right).
70
region was much smaller than the secondary electron image would suggest, due to a reaction layer surrounding the particle (Figure 9).
4. Discussion To understand the corrosion behavior of the alloys and the reason for the strong influence of Fe content, the mode of dissolution must be discerned . Some insight can be gained by study of the dependence of corrosion rate on immersion time shown in Figures 1 and 2. A notable feature of the curves is the rise in dissolution rate during the initial 5 to 10 hours. This behavior is most evident in specimens Band C, which contained the highest Fe concentrations, 0.009 and 0.020wt%, respectively. Since pitting was visually observed on these specimens, the initial rise is probably a result of pit initiation and formation at certain favorable sites on the surface. Once the pits are well developed , i.e. increasing in depth rather than diameter, the corrosion rate will be limited by ion transfer into and out of the pit and will stabi lize. The increase in rate for Band C may also be related to the delayed activation or exposure of pitting sites. Once all the available sites are active, the stable corrosion rate for the specimen will depend largely on their number and size. If one assumes that the principal active sites are due to Fe particles, then for particles above a critical size, there should be a correspondence of corrosion rate with the number of pitting sites. There is also an effect of particle size in that larger particles will provide more surface for the liberation of hydrogen gas and produce a larger local current density around the particle . Thus, the presence of either larger or more Fe particles in extrusion C accounts for its higher corrosion rate and corresponds with the wet chemical analysis. Unfortunately, because of the very low concentrations of Fe, only very few particles can be located on a specimen of manageable size. In addition, some pit-initiating particles are probably lost in the corrosion process once the pit has developed and undermined the particle. Simple calculation shows that Mg-15wt% AI, containing 0.009wt%Fe, will contain only 27 pure Fe particles of 10 !-tm diameter in a cross section of one square centimetre. Thus , insufficient numbers of particles were located for statistical characterization of particle size and number distributions.
Comparison of the actual weight loss corrosion rates with the EWLCR values from the hydrogen evolution tests showed good correspondence for all alloys (Figure 2), and demonstrated the validity of the technique. However, in all cases there was a consistent negative deviation from the true weight loss values of the corrosion rate determined by hydrogen evolution. This type of deviation has been noted previously in tests of this type and is generally attributed to the negative difference effect [6, 10-15]. In the present study, a large proportion of the weight loss discrepancy can be
FIGURE 9 (a) EDAX spectrum of Fe-rich particle with reaction zone in extrusion C. (b) Back-scattered electron im age of particle .
72
accounted for by assuming that the Mg17A112 particles are not dissolved. The presence of Mg17A112 in the corrosion product filtrate and the observation of barely visible shiny particles in the corrosion product (Figure 10) supports this assumption.
The effect of Fe content on the corrosion rate, plotted in Figure 3, with the results of Hanawalt for lOwt%AI shows that processing route also has a strong effect:!:. Hanawalt's alloys were cast in refractory molds, in which the alloys solidified at typically low rates. The comparison thus indicates that a higher solidification rate decreases the sensitivity to Fe content. This improvement has also been observed in other RSP Mg alloys [22, 23], in die cast Mg-AI alloys [1, 5], for RSP AI-base alloys [24], and has been generally explained in terms of reduced segregation. Rapid solidification may reduce both the size of cathodic second phases, and their volume fraction, by solubility extension. Counter to this effect, for Mg-AI alloys, is the presence of AI, which increases the amount of Fe-bearing phase by compound formation.
FIGURE 10 Cross-section of metal/corrosion product interface in extrusion B showing shiny Mg17A11 2 particles in corrosion product.
:j: The comparison between 10 and 15wt%AI is valid in this case because Hanawalt clearly showed that higher Al contents exaggerate the effect of Fe in mold-cast alloys. It is therefore, a conservative comparison.
73
The solubility of Fe in Mg is extremely limited and they form no compounds [25]. Therefore, the addition of aluminum to Fe-contaminated Mg would readily produce precipitation of AI-Fe compound(s) which have melting points well above alloy melt temperatures (assuming no ternary compounds). This would result in solid AI-Fe particles suspended in the melt to be incorporated into the Mg-AI powder particles resulting from atomization . Subsequent consolidation/extrusion of the powder would produce a fine dispersion of Mg17AI12 precipitates , as indicated by the X-ray results, in a Mg matrix , with occasional Fe-rich particles . Thus, RS of impure Mg-15wt%AI causes the Fe-rich particles to be in contact with Mg17AI12 precipitates encased by a Mg matrix . In contrast, conventional slow solidification rates will result in large dendrites of relatively pure Mg surrounded by interdendritic Mg/Mg17AI12 eutectic. When an Fe-rich particle is placed in the microstructure , the developing pit will likely incorporate dendritic regions enriched in Mg. These two scenarios are illustrated in Figure 11. There are three advantages, from a microgalvanic corrosion viewpoint, of the RSP microstructure : (1) the Mg17AI12 precipitates are passivated and apparently not dissolved to an appreciable extent by the solution. This creates a more tortuous path for ion transfer around the particles during matrix dissolution in a pit; (2) the precipitates are left imbedded in the corrosion product, possibly with some degree of interconnectivity, and may lend strength to the product which retards spalling; (3) a high density of the particles may alter the composition of the protective film which forms naturally in an aqueous environment to one which is more stable.
Conclusions 1. The aqueous corrosion rate of RSP Mg-15wt%AI alloy is markedly
augmented by minute quantities of iron, as shown earlier for mold and die cast alloys of this system .
2. Rapid solidification decreases the sensitivity of corrosion rate of Mg-AI alloys to iron impurity content. This is considered to result from the presence of passive precipitates in the microstructure which diminish ion transfer within pits , stabilize the passive surface film, and strengthen the corrosion product.
3. Filiform-like corrosion is observed, but is probably not a major contributing mode of corrosion in this case .
4. Magnesium hydroxide is the dominant corrosion product. Another compound is also produced by the RSP Mg-15wt%AI alloys which may be a form of hydrotalcite .
74
Corrosion Pit
Fe Particle
Solution
Corrosion Pit
Spalled ~ ____ Corrosion
• Product
(a)
Pits Containing Corrosion ________ Product and Passivated
Solution
(b)
FIGURE 11 Schematic corrosion pit formation in : (a) conventionally-solidified coarse microstructure with large Mg17AI1 2 particles not shown; (b) rapidly-solidified microstructure with fine distribution of Mg-AI precipitates.
75
Acknowledgements This work was the result of the 1989 F.M. Beckett Memorial Award for overseas summer study provided by the Electrochemical Society and was conducted at the University of Sheffield, Sheffield, England . Sincere appreciation is expressed to the Electrochemical Society for providing this opportunity, and to the faculty, staff and students at the University of Sheffield for their cooperation and guidance . In particular, the contributions of Elaine Leith, Robert Edyvean and David Rugg are gratefully acknowledged for their advice and prior research relating to this work . The RSP powder was atomized by Metalloys Ltd., as part of a study sponsored by the Aluminium Powder Co. Ltd ., Magnesium Elektron Ltd., Shell Research Ltd. and the U.K. Department of Trade and Industry .
References 1. J.E . HILLIS , Light Metal Age, June 1983 , pp. 25-29. 2. A. JOSHI and R.E . LEWIS, in Advances in Magnesium Alloys and Composites, eds . H.G.
Paris and W.H. Hunt, TMS, Warrendale , Pa., 1988, pp. 89-103 . 3. J .C. SCULLY, The Fundamentals of Corrosion, Pergamon Press, Oxford , 1975, pp . 174-
5. 4. M.G . FONTANA and N.D. GREENE, Corrosion Engineering, McGraw-Hill Book Co .,
1978, pp. 49. 5. H.P. GODDARD etal, THe Corrosion of Light Metals , Wiley and Sons , Inc., New York,
1967, pp. 259-311. 6. R. GLICKSMAN, 1. Electrochem. Soc., 1959, 106, pp. 83-88. 7. T.R. BECK and S.G. CHAN, f. Electrochem. Soc., 1983,130, pp. 1289-1296. 8. A GALLACIO et ai, U.S. Army Armament R&D Report AD-B075 280L, Dover, New
Jersey , June 1983. 9. L. WHITBY, Trans. Faraday Soc ., 1933 , 29, pp . 415-25.
10. P.F. KING , f. Electrochem. Soc ., 1966, 113, pp. 536-539. II. J.L. ROBINSON and P.F. KING , f. Electrochem. Soc. , 1961 , 108, 36-41. 12 . M.E . STRAUMANIS, f. Electrochem. Soc., 1958, 105, p . 284. 13. H. TUNOLD et ai, Corrosion Science, 1977, 17, pp. 353-366. 14. M.E. STRAUMANIS and B.K. BHATIA , f. Electrochem. Soc. , 1963, 110, pp. 357-360. 15 . G.L. MAKAR and J. KRUGER, in Advances in Magnesium Alloys and Composites, eds.
H.G. Paris and W.H. Hunt , TMS, 1988, pp. 105-121. 16. J .D. HANAWALT, C.E. NELSONandJ. PELOUBET, Trans. A.I.M.E., 1942, 147,pp.
273-299. 17 . D.S. AHMED, R.G.J EDYVEAN , H . JONES and C.M. SELLARS, proprietary
research funded by the U .K. Department of Trade and Industry Consortium at the School of Materials , University of Sheffield, 1989.
18. C.B. BALIGA , P. TSAKIROPOULIS andJ .F. WATTS , Internat. f . Rapid Solidification , 1989, 4, 231-250.
19. S. GUHA, M.Sc. Thesis , Thayer School of Engineering, Dartmouth College , Hanover , New Hampshire , October 1987.
20. F. HEHMANN, F. SOMMER , H. JONES and R.G .J . EDYVEAN, f. Mater. Sci , 1989 , 24, pp . 2369-2379.
21. E .S. LEITH, Ph .D. work in progress , University of Sheffield/Shell Research Ltd. , 1989. 22. D. RUGG , Ph. D. research in progess , School of Materials , University of Sheffield, 1988. 23. D.L. ALBRIGHT, in Advances in Magnesium Alloys and Composites, eds. H.G. Paris and
W.H. Hunt , TMS, 1988, pp. 57-75.
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24 . J. R. PICKENS, K.S. KUMAR and T.J. LANGAN, Proc. 33rd Sagamore Army Materials Research Conf., held at Burlington, Vermont, July 28-31 , 1986.
25 . W.G. MOFFATT, The Handbook of Binary Phase Diagrams , General Electric Co ., Schenectady , New York , 1981.
(Received 2nd February 1990 and accepted 2nd May 1990)
Specimen Preparation for TEM Studies of Rapidly Solidified Powders
N.J.E. Adkins, A.F. Norman and P. Tsakiropoulos
Department of Materials Science and Engineering, University of Surrey, Guildford, Surrey GU2 5XH
Abstract Rapidly solidified (RS) powders are either too large in cross section to be electron transparent or too small in extent to allow conventional TEM preparation methods. Three techniques (microtoming, nickel plating-powder sedimentation and copper plating-powder adhesion) are described which allow the preparation of electron transparent sections of non-ferrous alloy powders with little or no damage to the rapidly solidified microstructure . Typical results for all three techniques are presented and the benefits and shortcomings of each technique are discussed. The method of copper plating-powder adhesion allows a wide range of powder particle sizes to be plated simultaneously .
1. Introduction The microstructure of gas atomised powders can be complex due to the different heat extraction rates experienced by different size droplets, the different undercoolings achieved prior to nucleation and the distribution of active nucleants among the droplets [1]. Any study of the structure and phases produced in RS powders should consider the microstructure of both fine and larger particles. Such studies require detailed TEM investigations.
The specimen preparation method and the particle size of powder that can be examined under a transmission electron microscope (TEM) depend on the kind of powders and on the types of electron microscope available. In both cases, the first consideration is whether the powders are electron transparent at the available accelerating voltages. In this paper the term electron transparent is used to mean specimens that are sufficiently transparent to 100 to 200 kV electrons to allow bright field imaging . Powders produced by electrohydrodynamic atomisation (EHDA) are generally transparent to the 100 kV electrons used in conventional TEM analysis allowing the examination of the whole solidification microstructure [2] . However high pressure gas atomised (HPGA) powders (size range 1 to 200 ~m) or powders produced from pulverised RS ribbons «200""m) are
