Here is mine just in case you're interested.
1. Introduction
The aerospace industry has an ongoing requirement for strong, light weight materials in order to improve efficiency of aircraft and reduce the use of finite resources. New joining and fastening techniques can provide this, but need to be studied closely in order to ensure that they are reliable enough to use in safety-critical applications. A reduction in weight of the aircraft can also increase the payload it can carry, which improves efficiency. Aircraft are also being pushed beyond their design life span, so the long term corrosion properties of new alloys are very important to understand to keep the aircraft safe in service.
Friction Stir Welding (FSW) is a relatively new solid state joining process, which has successfully been used by NASA on the external fuel tanks for the Space Shuttle Endeavor [1]. There are regions in the weld that are subjected to high temperatures, which changes the microstructure of the alloy and hence the properties. Post weld heat treatments may be needed in order to restore the properties of the affected regions of the weld.
The alloy used for Endeavor’s thrust panels is a light weight aluminium-lithium alloy 2297. The 3rd generation, Al-Li 2050 is one, of these alloys are being increasingly used in advanced applications such as the space shuttle fuel tanks and thrust panels. The research on these alloys means that more applications are becoming available for them and Al-Li alloys are thought to be the next generation of alloys for use in the civil aerospace industry.
In friction stir welds, the nugget is subjected to high deformation and temperatures close to the melting point of the alloy [13] resulting in a re-crystallised microstructure. Depending on the processing parameters the cooling rate will vary, producing precipitates of different types and sizes in the matrix and on grain and sub-grain boundaries. In order to study the corrosion behaviour of such microstructures in a systematic manner, a sample alloy will have to be solution heat treated and then aged at different temperatures that are known to correlate with certain types of precipitates [6].
The aim of this project is to relate the corrosion behaviour of the alloy with certain type of precipitates, in different locations; and to determine the resistance to corrosion for each type of microstructure. The effect on the hardness of the alloy will also be determined.
2. Literature Review
2.1 Aluminium Alloys Containing Lithium
Lithium is the smallest metallic element and therefore the least dense. When it is alloyed with aluminium, each 1% weight of lithium added can reduce the density by up to 3% and increase the modulus by around 6% [2]. In the most recent generation of alloys, lithium is added to aluminium alloys also containing copper, magnesium, zirconium and silver [3]. Some examples of these alloys are AA 2195 and AA 2050. 2195 has the typical composition (Al–4Cu–1Li–0.4Mg–0.14Zr–0.4Ag wt%)[4] and 2050 has the typical composition of (Al-3.5Cu-0.8Li-0.36Mg-0.43Mn-0.09Zr-0.35Ag wt%) [5]. Al-Li 2195 is being used in the Space Shuttle external fuel tanks, as it has good low temperature mechanical properties [6, 7]. Al-Li 2050 is thought to be the future replacement for aerospace alloys such as AA 2024 (Al-4.5Cu-1.5Mg-0.6Mn wt% [8]) due to its weight saving benefits [5].
2.1.1 Al Alloys with Li and Cu
Adding copper not only increases the strength of the alloy, but it also reduces the solubility of lithium in aluminium. This leads to the formation of strengthening precipitates, such as T1, T2, δ and δ’ [3]. However, when copper is added to aluminium, the corrosion resistance increases if the copper is in solid solution, but when certain heat treatments are applied the copper can diffuse to grain boundaries, leading to inter-granular attack.
2.1.2 Effect of Various Elements in Al-Li-Cu-X
The addition of magnesium further reduces the solubility of lithium, even at high temperatures. Magnesium also helps strengthen the alloy by forming precipitates with copper and aluminium [3]. Adding zirconium refines the grain size and therefore improves the fatigue properties of the alloy. Silver and manganese can also be added to the alloy, as these elements assist with nucleation of precipitates [9].
2.1.3 Heat Treatment of Al-Li-Cu-X Alloys
A number of different precipitates can form in Al-Li-Cu-X alloys depending on the heat treatment applied (temperature and time). After solution heat treating the alloy to homogenise the structure, it is quenched to form a super-saturated solid solution. The alloy is then aged at various temperatures to nucleate and grow strengthening precipitates.
Precipitate Composition
GP zone Cu
δ AlLi
δ' Al3Li
T1 Al2CuLi
T2 Al6CuLi3
' Al2Cu
'' Al2Cu
Table 2.1.1 The composition of common precipitates found in Al-Li-Cu alloys [5].
It is known that T1 is the most effective precipitate for increasing the strength and fracture toughness of the alloy [10]. This is due to the plate shaped precipitates that are relatively large in comparison to the other precipitates, seen using a transmission electron microscope (TEM). T1 can form on sub-grain boundaries or within the matrix, but only grows very large when the habit plane is almost parallel to the sub-grain boundary [6].
δ' forms at much lower temperatures than the other precipitates and can nucleate within the matrix or on the interface of other precipitates [9]. This means that δ' is a very abundant precipitate and also improves strength due to its distribution and small size. δ' does not form on sub-grain boundaries [6].
' and precipitates form at higher temperatures compared to the above precipitates. They form both in the matrix and in the grain boundaries and contribute little to the hardness and strength of the alloy [11]. '' forms at lower temperatures held for a relatively long amount of time. If the temperature is held long enough for '' to continually grow, this can lead to an increase in hardness [6].
Fig. 2.1.1 TTP curves for Al-Li 2195 following solution heat treatment [6]
Figure 2.1.1 shows the Time-Temperature-Precipitation curves for an aluminium lithium alloy 2195. As seen in Figure 2.1.1(b), Chen and Bhat showed that δ' does not form at sub-grain boundaries. The heat treatment used by Chen and Bhat was to solution heat treat at 950F (510°C) for 1 hour, stretch the sample by 3% and apply various aging treatments to produce the TTP diagrams.
2.2 Friction Stir Weld Process
The friction stir weld (FSW) process was patented by TWI in 1991 [12] [[13]. It is a solid state joining process that uses a rotating tool piece to mix the butted sheets of metal together [14]. There are four regions within a FSW joint; the weld nugget, the thermo-mechanical affected zone (TMAZ), the heat affected zone (HAZ) and the unaffected material.
The weld nugget is part of the TMAZ, but is differentiated from the TMAZ as the metal has been re-crystallised [13]. The nugget undergoes high levels of deformation, but is then re-crystallised by the high temperatures experienced at the tool tip.
The TMAZ is next to the weld nugget and is a region of plastic deformation caused by the rotating tool piece. This region experiences elevated temperatures and the microstructure of the metal can change. Re-crystallisation does not occur in this region, because aluminium can undergo high levels of plastic strain [13]. This means there is a high density of dislocations in the TMAZ.
The HAZ has no deformation, but is subjected to a degree of thermal energy [12]. The size of the HAZ depends on how conductive the metal is, the weld parameters, and the size of the adjacent TMAZ. Temperatures can be high enough for certain precipitates to grow in a very localised area; and in the region near the TMAZ, the alloy can be over-aged leading to increased susceptibility to corrosion and decreased strength. The regions of a friction stir weld are shown in Figure 2.2.1.
Fig 2.2.1 Schematic diagram of friction stir weld zone microstructure [15].
A - Unaffected material
B - Heat affected zone (HAZ)
C - Thermo-mechanically affected zone (TMAZ)
D - Weld nugget (Part of thermo-mechanically affected zone).
2.3 Corrosion of Al-Cu-Li-X
2.3.1 Oxide Film
High purity aluminium is resistant to corrosion due to the presence of a thin passive oxide layer (~2.5 nm) [16]. This acts as an insulator to stop the aluminium reacting with the environment and blocks the cathodic reaction. The oxide layer thickness can change under different conditions; it can be removed by extreme conditions such as pH levels between 1 to 4 and 10 to 14. The oxide layer can be increased in thickness by raising the temperature, being in a damp environment or being placed in boiling water [17].
Alloying additions mean that surface defects are present in most of the aluminium alloys in service. These defects are more cathodically active than the pure aluminium, so localised corrosion can occur [18].
2.3.2 Pitting Corrosion
Pitting occurs on the surface where intermetallics and precipitates are present. These have a thinner passive layer compared to the aluminium oxide, which is more conductive and therefore cathodically active. These precipitates act as pit initiation sites. Pitting tends to occur in chloride-containing solutions: chloride ions are drawn into the pit to balance the positive charge on the metal ions within the pit and allow the anodic reaction to take place [18]. The hydrolysis of these metal ions leads to acidification, followed by dissolution of the metal being favoured over repassivation (oxide film formation) [18]. Pits can act as stress raisers that can initiate fatigue crack growth, causing failure in some components. Figure 2.3.1 shows how pitting occurs.
Fig 2.3.1 Systematic diagram of a pit [18].
2.3.3 Inter-granular Corrosion (IGC)
Copper is more noble than aluminium so when present in solid solution it can improve the corrosion resistance by blocking dissolution of the aluminium [17]. During certain heat treatments, the copper can form as precipitates on grain boundaries. The formation of the precipitates on grain boundaries draws the copper in from the matrix, forming a copper rich precipitate at the grain boundary.
The diffusion of copper to the grain boundary also creates a copper-depleted zone either side of the grain boundary [18]. This area is more susceptible to anodic dissolution as it has a lower corrosion potential. To this extent, a micro-galvanic couple is set up with the copper rich precipitate and the depleted zone dissolves preferentially [17]. Figure 2.3.2 shows the copper rich precipitates forming on a grain boundary, from this IGC can initiate.
The copper depleted zone can be part of the precipitate free zone (PFZ) as precipitates cannot initiate due to lack of vacancies for diffusion to take place. Inter-granular corrosion can be more serious as it can occur at a faster rate than pitting [17].
Fig 2.3.2 Grain boundary region in Al-Cu-X alloys [18].
2.3.4 Inter Sub-granular Corrosion (ISGC)
Some precipitates tend to form on high angle grain boundaries, known as sub-grain boundaries, as the energy required for them to nucleate is lower than within the matrix. With the possible presence of precipitates on sub-grain boundaries, inter sub-granular corrosion (ISGC) can occur. However, the susceptibility of the sub-grain boundary being a site for nucleation can vary and therefore some will have more precipitates present than others and hence the likelihood of corrosion will vary. The mechanism is similar to IGC, but corrosion is observed within the grain rather than at the grain boundary and shows as distinctive lines that highlight the sub-grain boundaries making the grain size seem smaller than expected, shown in Figure 2.3.3.
Fig 2.3.3 SEM of IGC, ISGC and matrix corrosion present in the TMAZ region of a friction stir welded Al-Li 2050 alloy [5].
2.3.5 Intra-granular Corrosion
Intra-granular corrosion (or matrix corrosion) is similar to IGC and ISGC, but the corrosion attack is of the matrix and not the sub-grain or grain boundaries. The mechanism is similar, as a precipitate within the matrix has an alloy depleted zone around it, reducing the corrosion potential and allowing anodic dissolution to take place. It is seen as darker dappled areas within the matrix of the alloy as seen in Figure 2.3.3.
2.3.6 Corrosion of Welds
The temperatures reached during the welding process can change the microstructure of the alloy considerably. The microstructural changes can be localised due to the intense heat involved. This can lead to poor corrosion resistance in the areas in and around weld joints. Certain precipitates can form in localised areas, which can change the corrosion resistance of that region.
It was found that corrosion rates were significantly increased in weld regions due to the precipitates forming microscopic galvanic couples. As well as this, pitting was found to be extensive in the weld regions of a 2195 alloy [2]. Conversely, Corral et al. showed that, for two alloys (2024 and 2195), the potentiodynamic polarization curves were very similar for the FSW regions and the base material. It was also shown that the corrosion product build-up was similar in both regions [8].
As discussed earlier, T1 is the main strengthening precipitate and it forms initially at defects (dislocations), sub-grain boundaries and grain boundaries. If the alloy has undergone pre-deformation, such as the T8 temper condition, then there will be more initiation sites for T1 to form at and therefore more precipitates will be present [19]. Referring to Figure 2.1.1b, T1 initiates quickly and over a wide range of temperatures. ' however, is slightly slower to initiate and forms at higher temperatures. In the weld nugget, T1 will have more chance to form than ' and is therefore the main precipitate in this region.
As T1 precipitates form, a copper and lithium depleted zone (PFZ) is formed adjacent to the sub-grain boundary, near where T1 forms. ISGC has been correlated to T1 and the PFZ due to the alternate anodic dissolution of T1 and the PFZ [19]. As T1 undergoes anodic dissolution, the reactive lithium atom is preferentially dissolved from the precipitate resulting in noble copper enrichment in the precipitate. De-alloying causes the T1 electric potential to become more positive compared to the PFZ and the galvanic couple is reversed and the PFZ undergoes anodic dissolution. This in turn exposes more of the T1 particle and the process is repeated [19].
Varying the heat treatment (i.e. weld parameters and post weld treatments) causes different sizes of precipitates to form in differing locations. This changes the weld's susceptibility to undergo numerous forms of corrosive attack. Post weld heat treatments can be used to improve the corrosion resistance for a weld region. However, these need to be selected carefully as over-aging can occur.
2.4 Summary
Little is known of the exact corrosion properties of various areas or friction stir welds. To simulate the zones of the weld, a systematic heat treatment of alloy samples is required. The weld nugget can be simulated by solution heat treating samples that can then be heat treated at various increasing temperatures. This will produce samples with varying types and amounts of precipitate formation leading to a difference in corrosion resistance and mechanical properties.
3. Experimental Method
3.1 Sample Preparation
A block was taken from a plate section of aluminium alloy 2050 in the T851 condition (solution heat treated, stress-relieved by stretching and then artificially aged [20]). The plate was supplied from Alcan - Centre de Recherches de Voreppe (CRV) in France. The plate had been friction stir welded on the top, but the block was taken from the bottom corner away from the HAZ. The block measured approximately 100mm wide, 35mm high and 70mm in depth. It was then sectioned into 10 equal sized samples that measured 10x35x70mm. At this point the rolling direction was noted and the samples were labelled S1 – 10 as seen in Figure 3.1.1.
Fig 3.1.1 Side view of the block sectioned into 10 samples.
S1 – 10 where used in this experiment as they will correspond to the heat treatment to simulate the weld nugget after solution heat treatment.
3.2 Heat Treatment
The ‘S’ (1 – 10) samples were solution heat treated at 525°C for 2 hours followed by water quenching. They were then aged at the following temperatures for 1 hour, 15 minutes.
Label Heat Treatment Temperature (°C)
S1 Un-treated
S2 100
S3 140
S4 230
S5 315
S6 400
S7 480
Table 3.2.1 Heat treatment temperatures for solution heat treated samples
All the samples were water quenched after the heat treatment. The temperatures were selected to be at points of interest in the TTP diagram of a similar Al-Li alloy (see Figure 3.2.1).
Fig 3.2.1 TTP curves for Al-Li 2195 adapted from Figure 2.1.1, showing the temperature in °F, and with the individual ‘S’ samples heat treatment temperatures in °C shown against the dwell time of 1 hour 15 minutes (dashed lines).
After the heat treatment, each sample was cut into 4 equal sub-samples that measured approximately 10x35x17mm and given the following label shown in Figure 3.2.2.
Fig 3.2.2 Top view of the sub-samples and the corresponding label.
S14 – S74 were then ground on the face shown in Figure 3.2.2 using silicon carbide paper and an ethanol based lubricant. They were then polished using a diamond suspension to a finish of 1 µm. After each grade change, the samples were rinsed in ethanol and finally cleaned in an ultrasonic bath.
3.3 Immersion
The samples were then coated with lacquer on all surfaces apart from the polished surface. The samples were then placed in a 0.1 M NaCl solution, which was naturally aerated, at room temperature and not stirred for 10 days. After 10 days, the samples were removed and rinsed in ethanol and air dried. Macro photos were then taken to display the amount of corrosion product and the severity of pitting. The samples were then cut through, at a point of interest, to reveal the cross section of the corroded surface and labelled ‘A’ and ‘B’. This produced further samples that were different in size, which were then ground and polished, as described above, in order that microscopy could be carried out on the surface of interest as shown in Figure 3.3.1.
Fig 3.3.1 Diagram of the corroded surface cut into ‘A’ and ‘B’ sections to show cross section of the corrosion morphology.
3.4 Microscopy
All SEM micrographs were taken using a JEOL 6060 scanning electron microscope using back scattered electrons at 15 kV.
3.5 Hardness
Macro-hardness testing was carried out using an Indentec Vickers Hardness testing machine with a load of 5 kg and a Vickers indenter. 5 indents were made on each sample 2 mm apart on the transverse cross section.
4. Results
4.1 Macro-Photos
Macro-photos were taken of the immersion tested samples before they were cut through, in order to determine the severity of pitting corrosion. A random 2 cm2 area was selected within the sample and the number of pits was counted for that area, to give a number of pits per cm2 for the sample. On some of the samples, the pitting was more densely distributed in a particular area, possibly because of increased localised precipitation formation. This is seen in the left hand end of S5 in Figure 4.1.1. The random area may not represent the full extent of pitting corrosion and some uncertainty in the results may be present.
Fig 4.1.1 Macro-photos of the immersed faces, clearly showing pitting corrosion occurring with some corrosion product on the surface of Al-Li 2050.
Sample Number of Pits / cm2
S1 16
S2 8
S3 1
S4 15
S5 28
S6 17
S7 21
Table 4.1.1 Number of pits per cm2 found on immersed samples.
S3 showed high resistance to pitting corrosion with only 1 pit on the whole surface. There is a drop in the number of pits as the temperature increases, but then increases with temperature to a peak at S5, dropping back briefly at S6 before a slight increase at the highest temperature.
4.2 Optical and SEM Microscopy
Fig 4.2.1 SEM micrographs of typical pits found in Al-Li 2050 after immersion testing. Samples were angled at 35° to show the top surface of the pit and the pit growth in the cross section of the sample.
Pit depth did not change significantly between samples however; the susceptibility to pitting corrosion did change as seen before in Figure 4.1.1.
Fig 4.2.2 Optical micrographs (left) of areas of corrosive attack on the cross section of the immersed Al-Li 2050 samples. High resolution SEM micrographs (right) of the boxed area, showing the corrosion morphology in more detail.
As the heat treatment temperature is increased, the susceptibility to IGC and ISGC is increased as seen in the SEM micrograph for S5. At lower temperatures there is an increased amount of intra-granular corrosion observed, compared with higher temperatures. This can be seen in S2 and S3 in Figure 4.2.2. The corrosion in S4 is a mixture of IGC, ISGC and intra-granular corrosion. The image for S6 shows a crack around a precipitate, inclusion or defect and only some matrix corrosion is seen around this feature. S1 and S7 are not included as there was nothing to note on the cross section of these samples.
4.3 Hardness
Macro-hardness measurements were taken on the cross section of the samples, 2 mm apart and roughly in the centre line of the sample.
Fig 4.3.1 Macro-hardness values against heat treatment temperature for Al-Li 2050 samples.
There is a drop in hardness until a sudden peak corresponding to S4, which is likely to have a high T1 precipitate content (Figure 3.2.1). There is another drop to the minimum hardness found in S6, before the hardness of S7 increases approximately to the value of S1.
5. Discussion
5.1 Corrosion
The solution heat treated alloy showed poor resistance to pitting corrosion (Figure 4.1.1). It was expected to have good corrosion resistance due to the homogenised microstructure. However, natural aging leads to the formation of GP zones [6], which takes noble copper out of solid solution and reduces the corrosion resistance.
S2 and S3 show good resistance to pitting possibly because of the dissolution of GP zones and θ’’ due to reversion [6]. This puts copper back into solid solution, mainly due to the GP zones dissolving, and general corrosion resistance from the noble copper is restored.
However, there was more intra-granular corrosion seen in S2 and S3 compared to the higher temperature samples. This relates to the matrix precipitates that form at slightly lower temperatures (mainly δ’) than precipitates that form at the sub-grain boundaries as seen in Figure 3.2.1 b. The presence of precipitates in the matrix can lead to intra-granular corrosion.
S4 and S5 will have a high proportion of the T1 Precipitate present (Figure 3.2.1). This precipitate has a high tendency to nucleate at sub-grain boundaries, and can grow very quickly from them [6]. This provides many pit initiation sites (Figure 4.1.1) and may be the cause of sub-grain boundary corrosion seen in S5, Figure 4.2.2. There is little intra-granular corrosion present in S5 due to the heat treatment temperature being un-favourable for precipitate initiation within the matrix (Figure 3.2.1 a). IGC and ISGC can occur due to the presence of the θ’ precipitate on sub-grain boundaries (Figure 3.2.1 b) There is a mixture of intra-granular corrosion, IGC and ISGC seen in S4 (Figure 4.2.2) due to precipitate formation being favourable in both the matrix and at the sub-grain boundaries (Figure 3.2.1).
There is a slight drop in the pits observed on the surface of S6 (Table 4.1.1). There is also little to no ISGC and IGC seen in Figure 4.2.2. This could be because the dissolution of T1 is putting copper back into solid solution, before the formation of θ or θ’ can occur.
At higher heat treatment temperatures the re-crystallisation and recovery process starts to take place, reversing any natural aging that may have occurred. However, before the precipitates can dissolve back into solid solution there is a pile-up of dislocations originated from residual stresses [6], which act as initiation sites and reduce the corrosion resistance. This explains the rise in pit formation seen in S7 (Figure 4.1.1)
5.2 Hardness
The drop in hardness from S1 to S3 could be explained by the reversion of GP zones and δ’. This means that there are less strengthening precipitates present to increase the hardness of the alloy.
The peak at S4 corresponds to the heat treatment that promotes the nucleation and growth of the T1 precipitate. This precipitate is the main strengthening phase because of its size, abundance due to the amount of favourable nucleation sites within the alloy and the rate at which it grows [6].
The hardness for S6 is unexpectedly low, as the presence of θ and θ’ in both the matrix and sub-grain boundaries should increase the hardness. This may not be true, because θ and θ’ have a tetragonal crystal structure, which has little effect on the hardness [11, 21]. If θ and θ’ are the only precipitates present in S6, then this could explain the sudden drop in hardness.
S7 corresponds to the heat treatment where re-crystallisation and recovery is starting to take place, so it is in a very similar state to S1 and shows a similar hardness as seen in Figure 4.3.1.
Fig 5.2.1 Hardness in Rockwell against temperature for an Al-Li 2195 alloy [6].
The hardness for a similar alloy is shown in Figure 5.2.1. The peak at around 450 °F (~230 °C) corresponds to the heat treatment that favours T1 nucleation and growth. This is similar in the 2050 alloy seen in Figure 4.3.1 at the S4 value for the hardness. There is also a drop in hardness for the first three treatment temperatures before the peak, which is similar to the behaviour of the Al-Li 2050 alloy. At higher temperatures, when the T1 precipitate is no longer present, the hardness drops off as seen also in Figure 4.3.1. The minimum hardness recorded for the 2195 alloy is when θ and θ’ are present (800 °F or 420 °C) which could explain the minimum hardness in the 2050 alloy. This can also be seen when the hardness values for Al-Li 2195 and 2050 are plotted together as shown in Figure 5.2.2
Fig 5.2.2 Hardness in Vickers for Al-Li 2050 and Rockwell for 2195 [6] against heat treatment temperature
6. Conclusion
6.1 Corrosion
A significant change in the corrosion attack was observed for different heat treatment temperatures. Natural ageing lead to an increase in corrosion attack in the solution heat treated sample (S1). There was a decrease in the corrosion observed for heat treatment temperatures between 100 °C and 140 °C due to the dissolution of GP zones because of reversion. The sample that was held at 140 °C for 1 hour (S3) showed the best corrosion resistance with only a single pit formed on the surface; no ISGC or IGC was found and there was only a small amount of intra-granular corrosion observed. Over-aging the alloy lead to an increase in the formation of pits, as well as more ISGC and IGC being observed.
6.2 Hardness
T1 was found to be the main strengthening precipitate because the sample held at 230 °C for 1 hour showed the highest hardness. However, the sample that showed the best corrosion resistance also showed very low hardness meaning that a compromise of the heat treatment is necessary to achieve good corrosion resistance, but maintaining the mechanical properties at the same time. The formation of θ and θ’ had little effect on the hardness of the alloy, and the samples that contained these precipitates showed the lowest values for hardness.
A more detailed study of the effect of heat treatment on the properties of this alloy is required to find the most suitable compromise for specific applications. A suitable heat treatment that maintains mechanical properties and corrosion resistance for this alloy would be between 140 °C and 230 °C for more than 1 hour.
7. References
1. Roy, S. Space Shuttle Endeavor's External Tank to Fly With Improved Welds, Lighter Weight Thrust Panels During Feb. 7 Launch. 2010 [cited 28/03/10]; Available from:
http://www.nasa.gov/centers/marshall/ne ... 0-010.html.
2. Walsh, D.W., The Effect of Fabrication on Corrosion in Aluminium 2195. 1994, California Polytechnic State University. p. 1 - 7.
3. Niskanen, P., T.H. Sanders, J.G. Rinker, and M. Marek, CORROSION OF ALUMINUM-ALLOYS CONTAINING LITHIUM. Corrosion Science, 1982. 22(4): p. 283-&.
4. Zhao, X., Corrosion behavior of 2195 and 1420 A1-Li alloys in neutral 3.5% NaCl solution under tensile stress. Trans. Nonferrous Met. SOC. China 16 1171-1177, 2006.
5. Kraft, S., Corrosion Investigations on a Multipass Friction Stir Weld Joint in Al-Li-Cu Alloy 2050-T8. 2008, Thesis, University of Birmingham. p. 1-40.
6. Chen, P.S. and B.N. Bhat, Time-Temperature-Precipitation Behavior in AI-Li Alloy 2195. 2002, IIT Research Institute, Huntsville, AL. p. 1 - 20.
7. Joshi, A., The new generation Aluminium Lithium Alloys. Metal Web News, 2008.
8. Corral, J., E.A. Trillo, Y. Li, and L.E. Murr, Corrosion of friction-stir welded aluminum alloys 2024 and 2195. Journal of Materials Science Letters, 2000. 19(23): p. 2117-2122.
9. Wang, Z.M. and R.N. Shenoy, Microstructural Characterization of Aluminum-Lithium Alloys 1460 and 2195. 1998, National Aeronautics and Space Administration, Langley Research Center. p. 1 - 46.
10. Li, H.Y., Y. Tang, Z.D. Zeng, Z.Q. Zheng, and F. Zheng, Effect of ageing time on strength and microstructures of an Al-Cu-Li-Zn-Mg-Mn-Zr alloy. Materials Science and Engineering a-Structural Materials Properties Microstructure and Processing, 2008. 498(1-2): p. 314-320.
11. Shukla, A.K. and W.A. Baeslack Iii, Study of microstructural evolution in friction-stir welded thin-sheet Al-Cu-Li alloy using transmission-electron microscopy. Scripta Materialia, 2007. 56(6): p. 513-516.
12.
http://www.twi.co.uk. [cited [28/02/10].
13. Zhi-hong, F., H. Di-qiu, and W. Hong, Friction stir welding of aluminum alloys. Journal of Wuhan University of Technology--Materials Science Edition, 2004. 19(1): p. 61-64.
14. Mishra, R.S. and Z.Y. Ma, Friction stir welding and processing. Materials Science & Engineering R-Reports, 2005. 50(1-2): p. 1-78.
15.
http://materialteknologi.hig.no/Lettvek ... r/fsw4.gif. [cited [22/03/10].
16. Polmear, I.J., Light Alloys : Metallurgy of the Light Metals. Third ed. 1993: Arnold.
17.
http://aluminium.matter.org.uk. [cited [01/03/10].
18. Davenport, A.J., Corrosion of Metals in Wet Environments: Corrosion of Aluminium Alloys, in Corrosion Lectures. 2009, University of Birmingham.
19. Li, J.F., Z.Q. Zheng, N. Jiang, and S.C. Li, Study on localized corrosion mechanism of 2195 Al-Li alloy in 4.0% NaCl solution (pH 6.5) using a three-electrode coupling system. Materials and Corrosion, 2005. 56(3): p. 192-196.
20.
http://www.engineersedge.com/aluminum_tempers.htm. [cited [24/03/10].
21. Cornell, R. and H.K.D.H. Bhadeshia. [cited 27/03/10]; Available from:
http://www.msm.cam.ac.uk/phasetrans/abstracts/M24.html.
FAQ
Posted: Wed May 19, 2010 8:34 pm 
