Six different alloys were cast in this experiment: See nomenclature (appendix C) for explanation of the different names. Three specimens of each alloy were cast to obtain more accurate results. Also, three different solidification rates were used for each alloy to simulate pressure die casting, gravity die casting and sand casting. Another batch was made for the heat treated bars; these were not studied through a microscope but were used in the tensile strength tests. The following chart shows the composition of each alloy. See appendix D for a complete chart of all elements.
Only Fe and Mn are varied, and as stated before the purpose of this thesis was to investigate the effects of Mn on Fe. Mn is gradually increased and Fe is kept constant after A16MB.
The alloys were produced by adding Si, Fe and Mn to the Aluminum melt. When it comes to Fe and Mn they were added in powder form with a purity of 75% respectively; the remaining 25% is some kind of flux media. The resistance furnace was set to 750°C for melting of the pure Aluminum and the alloying elements were preheated in a preheating furnace to 200°C. The melt was always skimmed before any addition of an element. The materials were added in following order: Si, Fe, Mn and last Sr.
In order to determine the composition of the alloy that had been produced, coins, one at the start and another at the end of the casting, were cast and analysed. The coins were turned using a lathe, which gave them a plane surface so they could be used for analysis. When performing optical emission spectrometry the surface of the sample is placed on a plate that is protected from oxides with help of protection gas argon, Ar. Thereafter a bit of the surface is burned and the amount of the different elements is presented.
Bars were cast for the gradient solidification experiments. With gradient solidification techniques it is possible to achieve a cast free from oxide films and porosity defects. In this work a resistance-heated furnace with an electrically driven elevator was used, see figure 2. This equipment offers the possibility to control the solidification rate and direction of the remelted metal. When re-melting the bars, a bar was inserted into the furnace, enclosed within a graphite coated steel tube. The metal was remelted in a protecting atmosphere (Ar), which prevents oxidation. The furnace moved with a constant velocity, v, and the bar was continuously cooled in water. The water was sprayed directly on to the steel tube. In this experiment 108 bars were cast. Three different velocities, v, were used, 0.03 mm/s, 0.3 mm/s and 3 mm/s. At every velocity, three bars were made. Fifty four of the bars were heated treated before tensile testing.
The procedure of heat treatment that was conducted follows the sequences presented below:
1. Solution heat treatment for 6h at 520 degrees.
2. After solution heat treatment samples were quenched in water of 60 C.
3. The final step is the warm aging of the samples for 8h at 165 C.
After the heat treatment procedure the bars were prepared for tensile testing.
Mechanical property tests
Room temperature tensile testing
All bars that were produced were used to determine the room temperature tensile properties of the alloy. The kind of tensile specimen that was prepared for tensile testing is schematically shown below both before and after the preparation, see figure
3. The as-cast specimens were tested at least one week after they were produced, while the heat treated ones could be tested directly after the heat treatment process. The tensile tests were conducted using a tensile test device called Lloyd EZ50, see figure 4 below.
Turn was done to avoid breakage tip, which can easily arise when using square-shaped samples. Afterwards the turning operation sample had a waistline diameter of 7 mm and waistline length of 60 mm including radii of 5 mm at the end of the waistline length, on both sides. The reason for these measurements depends largely on accommodation to tensile testing equipment EZ50.
Tension was accomplished with increasing tension load until it resulted in fracture; the elongation rate was 0.5 mm/minute. Strain was measured using an axial extensometer with a gage length of 25 mm, see figure 4 below.
The data in form of load and displacement was monitored and also stored and analyzed using Matlab to obtain ultimate tensile strength, yield strength, modulus of elasticity, elongation etc.
When the tensile tests were done, each of the specimens was prepared for microscopic examination. The preparation consisted of sectioning and mounting the specimens in a plastic medium to form a cylindrical piece, see fig 5, followed by grinding and polishing procedures of the specimens.
The phases have been identified by the color and morphology and compared to literature. No Energy Dispersive Spectrometry (EDS) analysis has been performed.
The influence of Mn and cooling rate on the microstructure
The following pictures were analyzed in an optical microscope in order to study the SDAS of the different alloys. Three pictures of each alloy are taken; one for each solidification rate, see figure 6. They can be seen as going from fine to coarse in that order. As observed, the solidification rate seems to be the most important factor. They all show the same SDAS for a given solidification rate, independent of the amount of Mn and Fe As demonstrated in figure 7, the solidification rate clearly has the greatest impact on the coarseness of the microstructure. The different levels of Mn did not seem to influence the dendrite cell size, neither did the Fe, see figure 7 (a-i).
Figure 8 shows intermetallic phases of the alloys at different cooling rates. The pictures illustrate how Mn affects the size of the needles.
The Fe needles are not visible for all cooling rates, the fastest speed did not allow enough time for needles to grow. Also the medium cooling rate is shown in 100 x magnifications for figure 8a and 8c. Only alloy F showed needles at SDAS 10Um. The location of the intermetallic regions is important to observe when evaluating these samples. This will be discussed later.
Each test from B-C shows that the needle length increases with slower cooling. Alloy D with 0.15% Fe and 0.15% Mn had both needles and Chinese script in the sample with slowest cooling, Fe needles dominated however. Notice in the picture above the Chinese sign is much larger than the needles.
Figure 9 shows alloy E and F. The alpha phases dominated these samples, and are compared to the sludge that also forms.
The graph (fig.10) shows an interesting trend in the last three alloys (MD-MF) at medium and slow cooling. The iron needles were found to increase in length with high Mn. The fastest solidification rate had needles visible in only one alloy and therefore was not plotted. This observation can perhaps be explained by the timing of the formation of the needles as was discussed before in the literature survey. The location of the needles gives a clue as to when they were formed. Intermetallic phases found in the dendrites must have been formed before the dendrites solidified (fig.11b), which is harmful as the intermetallic phases will grow much bigger then. It seems that higher Mn levels cause earlier formation of the intermetallic phases.
Figure 11(a-b) shows the alloy MF first at 100x Mg. and then at 5x Mg. for the slowest cooling. Even though there is 1% in this alloy the needles were still present, meaning that they are not completely suppressed. Figure 11(b) shows an example of what was discussed earlier about the location of the Chinese scripts with increased Mn levels.
The influence of Mn and cooling rate on the mechanical properties
The following graphs illustrate the influence of Mn on the mechanical properties of the different alloys (see appendix E for values used). The graphs are shown in as-cast and heat treated form, next to each other. An interesting notation is that Mn only improves the properties overall for UTS, YS and K, but not significantly. Note also that alloy A containing 0.1%Fe and alloy B containing 0.3% Fe do not differ much from each other, indicating that 0.3% iron may be an acceptable level for casting, where Mn is not needed for modification.
The effect of Mn on Elongation
Figure 12(a) shows how elongation is influenced by Mn content for as-cast test bars. The effect of Mn is roughly the same as heat-treated but with lower elongation values for all the samples.
Figure 12(b) shows the relation between elongation and Mn content for heat-treated test bars. At faster solidification speeds (SDAS 10 Um) the elongation is highest and the effect of Mn on Fe improves the elongation first at a Mn content exceeding 0.6 %. Overall however, Mn decreases elongation. The first two alloys with no Mn have a higher elongation than the other alloys.
The effect of Mn on the Strength Coefficient (K)
As cast samples
Figure 13(a) shows the relation between the material’s strength coefficient and Mn content for as-cast test bars. In this case the solidification time has considerable effect on K. Heat treatment lowered K for the “SDAS 10 Um” samples and raised it for the “SDAS 48 Um” samples.
Heat treated samples
Figure 13(b) shows the relation between the material’s strength coefficient and Mn content for heat-treated test bars. It looks like an optimal Mn content is between 0.4 and 0.8 percent. Mn shows a stronger effect on the as-cast samples.
The effect of Mn on UTS
Figure 14(a) shows the effect of Mn on the ultimate tensile strength for as-cast bars. The graph shows an improvement in UTS with the increase of Mn for the “SDAS 10 Um” and “SDAS 18 Um” test bars. At 0.6% Mn the curve levels out and no further improvement is seen. For the “SDAS 48 Um” bars, Mn decreases UTS until 1% is added, when it improves slightly. Overall UTS is not raised.
The effect of Mn on Yield Strength
Figure 15(a) shows the optimal effect of Mn on yield strength (YS) for as-cast test bars. The graph shows a clear overall improvement on all test bars when Mn is increased.
The effect of Mn on the strain hardening exponent
Figure 16(a) shows the effect of Mn on the strain hardening exponent of as-cast test bars. The greater SDAS have lower (n).
Fe in cast aluminum alloys has a detrimental effect on mechanical properties and is preferably kept as low as possible. In the absence of modifiers the iron forms large plate-like structures which hinder interdendritic feeding and cause porosity due to shrinkage. Also, cracks tend to start and spread through them. Mn is added to change the morphology of these intermetallic phases to less harmful ones. It is also added in some cases to reduce sticking to the die. Excessive amounts of Mn will cause sludge problems however.
The cooling rate and Mn content has the biggest impact on the formation of intermetallic phases. Fe needles were hard to discern and mostly were not observed at the highest cooling rate, even with 100x magnification. This is because they don’t have time to grow during solidification. As the cooling rate decreases, more Fe needles are seen. The same observation was made about the Chinese script.
The cooling rate also affects the porosity level. The highest solidification rate showed smaller pores that were dispersed evenly throughout the material. There are less Fe needles in these alloys that would block feeding and also the finer structure allows a better flow of the melt which in turn does not allow for much pore formation. As cooling rate decreased the pores became larger and appeared in local concentrations. According to Campbell  lengthy cooling rates often result in shrinkage porosity which could be an explanation to why much larger pores were encountered in the slowly (0.03mm/s) cooled series. It is also possible that the large pores are formed when bifilms in the melt expand during cooling .
With the slowest cooling rate, there were many factors that could have caused differences in mechanical properties so it was more difficult to see the direct effect of Mn. Factors such as SDAS and pores caused premature breaks in many of the test bars.
The first Chinese script phases were found in the MD test bars containing 0.3% Mn; and then only in the test bars cast at the slowest cooling rate. At values of 0.6% Mn and above the Chinese scripts were predominant. However, the harmful β-phase was never completely suppressed. Even in the MF test bars containing 1% Mn some Fe needles were still found; at a Mn/Fe ratio of 3. This is in contrast to an article written by A. G. Prigunova on the subject. He found that when the Mn/Fe ratio is greater than 1.2 for hypoeutectic alloys, only the α-phase forms . These observations were made for alloys containing 0.7-1.3% Fe however, which may possibly account for the difference.
An increase in the size of needles and Chinese scripts was observed with greater Mn. Higher Mn resulted in a reduced amount of Fe needles but larger ones. Not enough values were used to make this observation statistically reliable however. There was also a greater fraction of intermetallic phases with higher Mn levels. Elongation decreased with the addition of Mn, which can be seen as a direct result of these higher fractions and larger needles. It is not clear why the needles are bigger with higher Mn and this was only found to be true on the “SDAS-18Um” and ”SDAS-48Um” test bars (see figure 10), but it appears that too much Mn might counter the original purpose. This might not be so if all β-phases were suppressed. Perhaps the increases in Mn lead to an early formation of the metallic phase, as was discussed by Taylor . When the intermetallic phases nucleate before the primary phase has solidified they are given more time and room to grow. Note that the Chinese scripts are only seen in the primary phases on the samples with high levels of Mn.
While ductility decreases with higher fractions of intermetallic phases, ultimate tensile strength and yield strength improve. It is clear then that Mn does not improve ductility, as it creates higher fractions of intermetallic phases. This is in accordance to the article written by Rheinfelden, where they state that Mn will lower ductility when its content exceeds 0.2% .
The influence of Heat Treatment on mechanical properties
Heat treatment decreased the UTS, YS and K while it improved elongation and n. Heat treatment does not change the effect of Mn on the mechanical properties except for K. The as-cast graph shows an increase in K with higher levels of Mn whereas the heat treated graph stays level. A study made by Rheinfelden on an alloy called Castasil-37 showed that heat treatment does not change the effects of Mn on the mechanical properties. This is true for our samples also.
This research verifies observations made by Carlos & Nestor in their thesis on pore concentration and appearance at different solidification rates . They found that when Fe content was 0.3 % pore fraction decreased to 2 %, when adding 0.7 % Mn. This gives some support that Mn may decrease the pore fraction .When adding 0.3 % Mn at the same conditions as Carlos & Nestor the pore fraction was about 4 %. Twice the amount of pore fraction was observed with half the Mn content. One author in an investigation about engine block casting suggests a lot more strontium and low quantity of Mn in order to reduce micro porosity . Mn may reduce shrinkage pores as well [11, 16].
A greater fraction of pores was found in the medium cooling rate than in the slowest cooling. However the largest pores were found in the slowest cooling (see appendix F). This indicates that cooling has the biggest impact on the size of the pores and Mn has the biggest impact on the fraction of pores. Perhaps a larger fraction is found in the medium cooling rate because the modification process takes time. More time allows a full modification. This would be true up to a point; for the highest cooling rate the Fe needles are few and small and do not need to be modified. This is supported by Campbell in his theory of how pores are not able to grow in high pressure die castings due to the small size of the SDAS .
The effect of porosity on mechanical properties was seen during the tensile tests. Approximately 40% of the test bars had pores on the fracture surface and can be assumed to have caused the break. Most of these test bars had solidified slowly and the pores were large. Premature breaks in pore areas caused a decrease in elongation of those samples.
Table of Contents
1.1 PURPOSE AND GOALS
2 Literature Survey
2.2 CAST ALUMINUM ALLOYS
2.3 ALUMINUM SILICON ALLOYS
2.4 PHASES OF AL-SI ALLOYS CONTAINING IRON AND MANGANESE
2.5 MANGANESE IN CAST ALUMINUM ALLOYS
2.6 MANGANESE AND PRE-HEAT TREATMENT
2.7 IRON IN CAST ALUMINUM ALLOYS
2.8 PORES IN ALUMINUM ALLOYS
2.9 STRONTIUM IN CAST ALUMINUM ALLOYS
2.10 HEAT TREATMENT
3 Experimental Technique
3.1 ALLOY CHART
3.2 GRADIENT SOLIDIFICATION
3.3 HEAT TREATMENT
3.4 MECHANICAL PROPERTY TESTS
3.5 MICROSCOPIC EXAMINATION
4.1 THE INFLUENCE OF MN AND COOLING RATE ON THE MICROSTRUCTURE
4.2 THE INFLUENCE OF MN AND COOLING RATE ON THE MECHANICAL PROPERTIES
5.1 COOLING RATE
5.2 INTERMETALLIC PHASES
5.3 HEAT TREATMENT
5.4 PORE FRACTION
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