Recent progress in magnesium–lithium alloys

Pressure die casting

In pressure die casting, as the melt temperature of Mg–Li base alloys is usually about 600–700°C, the cold chamber die casting method is always used. In this process, the pressure is an important parameter. Regener et al. ¹ investigated the effects of two levels of pressure (I and II were used to identify the pressure values, whereby II is larger than I) on the microstructure and mechanical properties of Mg–10·7Li–0·6Al. The higher pressure results in the finer microstructure of the alloy (Fig. 1). The strength and elongation of the alloy are also increased by an increase of pressure (as shown in Table 2).

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1.

Effect of pressure on the microstructure of pressure die cast Mg–10·7Li–0·6Al

Semi-solid forming

There are several methods for semi-solid forming, such as mechanical stirring, electromagnetic stirring, and strain induced melt activation (SIMA). Among these methods, SIMA is the most commonly used in the preparation of magnesium alloys. In the SIMA process, the strain and temperature are the two main parameters affecting the microstructure and properties of alloys. Figures 2 and 3 show the effects of the strain and temperature on the microstructure of Mg–Li–Al–Ca alloys. ² A proper temperature produces a fine microstructure; too high or too low temperature makes the microstructure coarser. Higher strain results in finer microstructure. Through semi-solid forming, the strength of Mg–Li–Al alloys can be improved. Figure 4 shows the comparison in strength of the alloys fabricated by traditional casting and semi-solid forming. It is obvious that the strengths of the semi-solid formed Mg–Li–Al alloys are higher than those of as-cast alloys.

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2.

Microstructural changes of 15% strain-introduced Mg–Li–Al–2%Ca alloys under different semi-solid forming temperature

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3.

Microstructural changes of a Mg–9Li–3Al–2Ca and b Mg–14Li–3Al–2Ca alloys introduced various amount of strain. Specimens were heated at semi-solid temperatures of a 547°C and b 549°C

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4.

Strength of as-cast and semi-solid press-formed Mg–Li–Al alloys

Rapid solidification

In the preparation of Mg–Li base alloys, rapid solidification technologies include: twin roller quenching, melt spinning and spray casting.

Figure 5 is the schematic of the twin roller quenching method. In this process, the charge is melted in a tubular crucible. Then, the melt is shrouded by argon gas and expelled from the crucible. The melt coming from the tubular crucible is rolled and quenched by two rotating copper rollers. In this way, ribbon alloys are obtained. Meschter and O’Neal studied the microstructures of the alloys of Mg–9Li, Mg–9Li–1Si and Mg–9Li–1Ce prepared with the twin roller quenching process and the conventional process. Results showed that the twin roller quenching produces microstructural refinement by a factor of 10–30 as compared to conventional processing. ³

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5.

Schematic drawing of twin roller quenching process

Figure 6 is the melt-spinning apparatus. In this process, the melt was also shrouded by argon gas and expelled from the tubular crucible. The melt coming from the tubular crucible was quenched by a high speed copper wheel and a solidified ribbon sample was obtained. Matsuda et al. ⁴ used this process to prepare Mg–13Li–4Si–1Ag specimens with a thickness of 50–100 μm. The specimens possessed homogeneous composition, a fine and uniform distribution of the intermetallic phase (Mg₂Si, 2–3 nm), and a highly fine-grained matrix.

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6.

Melt-spinning apparatus

During the spray casting process, melt is sprayed in a nozzle by an atomiser and consolidated on a substrate disc (Fig. 7). Figure 8 (Ref. 5) shows the comparison of the microstructures of Mg–Li alloy under two casting conditions (chilled casting and spray casting). The precipitated phase in the alloy is needle shape in the chilled casting specimen, while in the spray casting specimen, the precipitated phase is particle shape and distributed in the matrix more homogeneously. The size of the precipitates is also refined. Responding to the change of microstructure, the hardness of spray casting specimens is higher than that of chilled casting specimens by 20%. ⁵

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7.

Schematic drawing of spray casting process

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8.

Comparison of Mg–Li alloy microstructures under different casting process

Formability

In Mg–Li alloys, the BCC β phase with a preferred orientation can show anisotropic behaviour, but this is less pronounced when compared with a hexagonal close packed (hcp) α phase with a similar degree of texture. The asymmetry of the hcp α phase will exacerbate the effects of texture.

In Mg–Li–Al alloys, the hexagonal magnesium-rich solid solution α phase exhibits a comparatively high flow stress and low formability, whereas the BCC lithium crystal β phase is soft and ductile. The intermetallic AlLi phase (ordered B2 structure) is macroscopically not deformable and increases the flow stress by dispersion strengthening. A higher volume fraction of the AlLi phase causes microcracks during deformation. The ductility increases by increasing amounts of the eutectic constituents (α phase plus β phase). The optimal combination of volume fractions of the α, β and AlLi phases enables an increase in strength with sufficient ductility. ²⁹

Wu et al. ^30,31 compared the mechanical and anisotropic behaviour between LZ60 and LZ90. With a low average n value of approximately 0·159, LZ60 did not exhibit good ability to resist plastic instability during tension. The r value of it is not large enough to give good drawability and the large negative value of Δr results in serious ear formation during the drawing process. Therefore, cold-rolled LZ60 thin sheet did not exhibit good stretchability and drawability at room temperature. As for LZ90 alloy, it presents excellent ductility even at room temperature. Thus, the formability is better than that of LZ60 alloy. Although LZ90 exhibited excellent ductility at room temperature, the strength level is somewhat inferior to that of LZ60 alloy, and the stretchability is not much superior to LZ60 because of its rather small n value. The value of the LZ90 alloy is smaller than that of the LZ60 alloy, but the higher elongation resulted in moderate drawability. The large negative values of Δr for the cold-rolled LZ90 thin sheet would also cause severe ear formation during the drawing process at room temperature. The increase of values and decrease of Δr values with increasing temperature revealed that the drawability of LZ90 alloy could be enhanced by forming the sheet at elevated temperatures. The forming limit diagram indicated that stretchability of the LZ90 alloy could be significantly improved by deforming the sheet at elevated temperatures, as shown in Fig. 15. ³²

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15.

Forming limit diagrams of LZ90 sheet at various temperatures

Trojanova et al. ³³ studied the compression mechanical properties of three alloys Mg–4Li, LA43 and LA45 at temperatures between 25 and 200°C. Compared with LA43 and LA45, Mg–4Li alloy depends significantly on the testing temperature. Above a certain temperature, softening processes influence the course of the work-hardening curve. Increasing the activity of softening processes with increasing temperature causes differences in the deformation behaviour. Cross-slip and climb of dislocations may be responsible for the observed softening. As for LA43 and LA45, dynamic strain aging causes the post-stress relaxation (SR) effect the flow stress at the beginning of plastic deformation after SR is higher than that at the beginning of the SR. The dynamic strain aging is because of the migration of dissolved Li and Al atoms towards dislocations. ³⁴

Takuda et al. ^35,36 investigated the deformation behaviour of Mg–8·5Li–1Zn alloy sheet (a thickness of 0·6 mm after warm rolling with a total reduction ratio of 85%) at a different strain rate. Uniaxial tension tests were carried out at room temperature for various strain rates between 1·4×10⁻⁵ and 8·3×10⁻² s⁻¹. Results show that the sheet has strain-rate sensitivity even at room temperature with a strain-rate sensitivity exponent m of 0·064. The sheet has sufficiently high formability at comparatively low strain rates. The maximum elongation reaches 100%. At high strain rates, the grains are divided into fine subgrains without any obvious textural change. At low strain rates, the orientation of the texture changes during the tension tests, especially in the β phase, and the grains are elongated in the tensile direction. In cylindrical deep drawing, bore expanding and Erichsen tests for the alloy show that the limiting drawing ratio, the bore-expanding ratio and the Erichsen value are 2·2, 0·8 and 9·0 mm, respectively. It is found that the good formability of the alloy is attributed to the high strain-rate sensitivity exponent.

Comparing the formabilities of Mg–6Li–1Zn, Mg–9·5Li–1Zn and Mg–12Li–1Zn, the ductility of Mg–6Li–1Zn is inferior to that of Mg–9·5Li–1Zn and Mg–12Li–1Zn. The relationships between the stress and the strain rate in log–log scale for the three alloys are shown in Fig. 16. The strain-rate sensitivity exponent m for Mg–6Li–1Zn is small and keeps a constant value of 0·02. The m values for Mg–9·5Li–1Zn and Mg–12Li–1Zn do not keep a constant value with an increase of strain rate. The m values for Mg–9·5Li–1Zn and Mg–12Li–1Zn are 0·12 and 0·14 at lower strain rates, and are 0·05 and 0·06 at higher strain rates, respectively. The m values for Mg–9·5Li–1Zn and Mg–12Li–1Zn are obviously higher than that for Mg–6Li–1Zn, which induces better ductility of Mg–9·5Li–1Zn and Mg–12Li–1Zn than that of Mg–6Li–1Zn. ³⁷

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16.

Relationships between true stress and strain rate in log–log scale for a strain of 0·1

Prasad et al. ^38,39 compared the processing map for hot working of hot-rolled Mg–11·5Li–1·5Al and Mg–11·5Li–1·5Al–0·15Zr. Hot-rolled Mg–11·5Li–1·5Al and Mg–11·5Li–1·5Al–0·15Zr alloys were hot compressed in a temperature range of 200–450°C and strain rate range of 0·001–100 s⁻¹. The map of Mg–11·5Li–1·5Al exhibits two domains, a peak efficiency of about 62% occurring at 450°C and 0·001 s⁻¹ and a peak efficiency of about 58% occurring at 300°C and 0·001 s⁻¹, respectively. In the domain at 450°C and 0·001 s⁻¹, intergranular cracking exists in the microstructure, which is accordingly not a safe domain for processing. In the domain at 300°C and 0·001 s⁻¹, the apparent activation energy for hot deformation is 110 kJ mol⁻¹ and a very high tensile ductility is obtained. Therefore, this domain is identified to represent superplastic deformation. ⁴⁰ For Mg–11·5Li–1·5Al–0·15Zr alloy, the map exhibits a single domain with a peak efficiency of 53% occurring at 350°C and 0·001 s⁻¹. This domain represents dynamic recrystallisation (DRX). Unlike in Zr-free alloy (Mg–11·5Li–1·5Al), which exhibits only superplasticity, the Al₃Zr particles in the alloy help in nucleating DRX and in refining the grains. The alloy exhibits flow instabilities at temperatures higher than 350°C and strain rates higher than 10 s⁻¹.

Superplasticity

Superplasticity is the ability of materials to exhibit large elongations, typically several hundred per cent, at elevated temperatures and slow strain rates. It has been reported that Mg–Li base alloys in the eutectic composition range, with their microstructure consisting of two phases of hexagonal α phase and centred cubic β phase, can produce superplasticity under appropriate thermo-mechanical processing and testing conditions. ^41–45 Superplasticity can improve the forming performance of magnesium–lithium alloys.

Cao et al. ⁴⁶ prepared Mg–7·83Li and Mg–8·42Li alloys. The alloys were hot rolled, warm rolled and cold rolled with four passes, two passes and two passes, respectively. The total reduction was 92·5% (from the thickness of 20 to 1·5 mm). Then, the alloy was annealed in a nitrate bath at 648 K for 30 min. An average grain size of 3·48 μm was obtained in the alloys. The superplasticity tensile curves for the alloys are shown in Fig. 17. The maximum superplasticities of 850 and 920% were obtained in Mg–7·83Li and Mg–8·42Li alloys at 573 K and at an initial strain rate of 1·67×10⁻³ s⁻¹ and at an initial strain rate of 5×10⁻⁴ s⁻¹, respectively. After the superplastic deformation at 573 K, the grains of Mg–7·83Li alloy grow. The grains maintain the characteristics of equiaxial ones after large deformation, indicating that grain boundary sliding (GBS) is the dominant deformation mechanism. Different atomic mobilities of Mg, Li, α and β phases and different Gibbs free energies of α and β phases contribute to the obvious grain growth. Lattice diffusion is the grain growth mechanism. The stress exponent n, grain size exponent p and activation energy Q were calculated experimentally to be 2·1–2·3, 1·9 and 107 kJ mol⁻¹, which are close to the theoretical 2, 2 and 112 kJ mol⁻¹, lattice diffusion activation energies, indicating that the dominant deformation mechanism is GBS accommodated by slip controlled by lattice diffusion.

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17.

The nominal stress versus nominal strain curves of superplastic deformation in Mg–7·83Li and Mg–8·42Li alloys

Mg–8·5Li alloy after high ratio (100∶1) extrusion at 623 K possesses good superplasticity (a maximum elongation of more than 500%) at high strain rate, 1×10⁻² s⁻¹, at 623 K. ⁴⁷ The strain-rate sensitivity exponent, m value, is about 0·40 at 623–673 K at strain rates ranging from 10⁻³ to 10⁻¹ s⁻¹. The apparent activation energy for the superplastic deformation at 573–673 K with m = 0·4 is calculated to be about 80 kJ mol⁻¹, revealing that the dominant deformation mechanism in this alloy is still GBS. Therefore, GBS and DRX during superplastic flow should be responsible for the superplastic characteristics observed in the alloys.

Mg–8·5Li–1Zn and Mg–8·5Li–3Zn alloys after high ratio (100∶1) extrusion at 623 K also possess high strain-rate superplasticity because the high ratio extrusion is effective to induce DRX in the alloys, leading to grain refinement. ⁴⁸ Mg–8·5Li–1Zn alloy shows a high elongation of about 400% at 623 K at a strain rate of 1·1×10⁻² s⁻¹, and the Mg–8·5Li–3Zn alloy demonstrates a maximum elongation of more than 540% at 623 K at a strain rate of 1·18×10⁻² s⁻¹. At the examined temperatures of 573–673 K within the m value around 0·48 range, the activation energy for superplastic flow in Mg–8·5Li–1Zn and Mg–8·5Li–3Zn alloys is about 86 and 79 kJ/mole, respectively, indicating that the dominant superplastic deformation mechanism in the two alloys is GBS controlled by grain boundary diffusion.

Mg–11Li–3Zn alloy after rolling a total reduction of 94%, with a final thickness of 1·2 mm, possesses quasi-superplasticity with a maximum elongation of 200% at a temperature of 573 K and strain rate of 1·67×10⁻² s⁻¹. ⁴⁹ During the superplastic tensile deformation, DRX occurs causing grain refinement, corresponding to an average grain size from 27 to 9 μm. The stress exponent is 4·4 and the activation energy is 112·6 kJ mol⁻¹. The superplastic deformation mechanism is dislocation climb controlled by lattice diffusion. A rolled sheet with a thickness of 0·6 mm and a composition of Mg–9·5%Li–1·0Zn exhibits a quasi-superplastic behaviour with a maximum elongation of 290% at 523 K with an initial strain rate of 1·0×10⁻⁴ s⁻¹. ⁵⁰ The experiments give a strain-rate sensitivity of ∼0·33 and an activation energy of ∼92 kJ mol⁻¹. Measurements of the shapes and sizes of the internal cavities produced during superplastic deformation reveal a transition with increasing strain from cavity growth by diffusion to growth controlled by plastic flow in the surrounding crystalline matrix.

Superplastic deformation has an influence on the mechanical properties of Mg–Li base alloys. Kaibyshev et al. investigated the influence of superplastic deformation on the alloys of IMV-2 and MA21. ^51,52 Superplastic deformation at high strain rates leads to high tensile strengths and yield points. The alloys possess highly stable mechanical properties because of the highly homogeneous structure after superplastic deformation.

Superplastic deformation also affects the microstructure of Mg–Li alloys. The as-cast and thermo-mechanically processed Mg–8Li alloy possesses a maximum elongation of about 300% under suitable superplastic conditions. During superplastic deformation, needle-like α phase changes to an equiaxial shape, and the m value of 0·2 at the initial period correspondingly increases during the deformation. Twinning is observed in the α phase at the initial stage of superplastic deformation. The α grains are found to grow during the deformation, while the refining of β grains takes place. ⁵³ The Mg–8·4Li alloy after rolling with a thickness of 2 mm possesses an elongation of 920% at 573 K at an initial strain rate of 5×10⁻⁴ s⁻¹ with a strain-rate sensitivity exponent of 0·64. ⁵⁴ The equiaxial fine grains and nearly 50/50 ratio of two phases α/β are responsible for the larger superplastic elongation. The average grain sizes at tensile specimen gauge length section before and after superplastic tension and at grip section are 7·5, 31·7 and 20 μm, respectively. This indicates that the superplastic deformation induces grain growth, and static grain growth occurs during the deformation at this temperature. The Mg–8Li alloy after being hot rolled to 3 mm thick at 623 K, and cold rolled to 1·5 mm thick with intermediate annealing temperature, exhibits a maximum elongation of 960% at a temperature of 573 K, with an initial strain rate of 5×10⁻⁴ s⁻¹. ⁵⁵ During superplastic deformation, DRX takes place. Grain growth and cavity nucleation at α phase grain boundaries can also be observed. It is postulated that the continuous introduction of lattice dislocations into the phase interfaces contributes to the superplasticity of the alloy.

Furui et al. ⁵⁶ compared the ductilities of Mg–8Li alloys processed by different thermo-mechanical procedures. Results show that Mg–8Li alloy possesses very limited ductility in the as-cast condition and only moderate ductility after rolling. However, excellent superplastic properties are achieved at a testing temperature of 473 K by extruding the cast alloy, and the superplastic properties are further improved by extruding and processing through two passes using equal-channel angular pressing (ECAP), as shown in Fig. 18. The Mg–8Li alloy processed by deformation combining the extrusion and ECAP possesses a maximum elongation of 970% at a temperature of 473 K, and with a strain rate of 1·0×10⁻⁴ s⁻¹. ⁵⁷ During superplastic deformation, the frequency of high angle boundaries, defined as above 15° misorientation, increases with increasing the elongation to failure in order to advance DRX. Cavities form at the boundaries between α and β phases from the elongations of about 600%. The area fraction of cavities in the cast+extrude+ECAP condition was smaller than in the cast+extrude condition. The measured strain-rate sensitivity was of the order of 0·6. The activation energy for superplastic flow was essentially equal to the value for grain boundary diffusion of Mg. ⁵⁸ The addition of Zr in Mg–8Li alloy induces the maximum elongation increase because of the finer grain structure before loading and suppression of dislocation processes by solute atoms of Zr, which results in both uniform deformation and DRX. ⁵⁹

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18.

Appearance of specimens processed by the different routes and subsequently tested in tension to failure at 473 K using the strain rate of 1·0×10⁻⁴ s⁻¹

Qu et al. ⁶⁰ prepared Mg–8Li alloy with the methods of casting and extrusion. The diameters of the specimen before and after extrusion were 55 and 15 mm, respectively. The extrusion temperature was 553 K. The grain size of the alloy was less than 10 μm. The superplastic property of the alloy was somewhat poor. Only at the temperature larger than 523 K and strain rate of less than 2×10⁻⁴ s⁻¹ the alloy exhibits superplasticity. A maximum elongation of 164·5% was obtained at a temperature of 563 K with an initial strain rate of 5×10⁻⁵ s⁻¹. The processes of formation and merging of cavitations are favourable for the superplastic behaviour of the alloy. When the size of cavitation is large enough, tensile failure occurs.

Yoshida et al. ^61,62 prepared Mg–Li–Zn alloys with repetitive ECAP processing, which possess low temperature superplasticity because of the refining of both the phases of α and β, and spheroidising of the α phase. The elongation of the alloy at 373 K, which is lower than T _m/2, reaches about 400% with a strain rate of 1×10⁻⁴ s⁻¹. The alloy has a relatively high strain-rate sensitivity exponent of about 0·4. During superplastic deformation, the DRX and particle precipitation occur in both α and β phases. An increase of grain boundary area through the fine recrystallised grains and precipitation of β phase on the recrystallised grain boundaries of the α phase contributes to GBS, resulting in the occurrence of superplasticity.

Mg–8Li–2Zn alloy processed by two-pass extrusion (553 K, from Φ106 to Φ50 mm, then from Φ50 mm to 30×3 mm) exhibits superplasticity at low temperature (an elongation of 120% at a temperature of 423 K) ⁶³ and at high strain rate (an elongation of 279% at an initial strain rate of 1×10⁻² s⁻¹). ⁶⁴ This alloy has the largest elongation of 758% at 563 K with a strain rate of 1·5×10⁻⁴ s⁻¹. ⁶⁵ During superplastic deformation, the strain-rate sensitivity is about 0·55 and the activation energy is about 89·4–99·24 kJ mol⁻¹, showing that the deformation is grain-boundary sliding controlled by grain-boundary diffusion. Coalescence and interlinkage of cavities is the reason for tensile failure.

Deformation texture

Deformation texture of alloys not only affects the subsequent deforming processes but also causes anisotropic mechanical properties in the alloys. Accordingly, texture in deformed alloys is a very important factor affecting properties of the alloys.

Mg based alloys are typically hcp structural materials. Compared with face-centred and body-centred cubic materials, hcp metals exhibit a wider variety of deformation textures. The distinct textures depend on the combined effects of c/a ratio and the fact that the plastic deformation occurs through different slip systems and/or twinning modes. ⁶⁶

The addition of Li and other alloying elements affects the texture of Mg based alloys. Agnew et al. ⁶⁷ investigated the compression textures of as-rolled Mg–(0–5)Li alloys. Results show that, when the addition of Li is less than 5 wt-%, there are only subtle differences in the compression textures of magnesium. However, the plane strain compression textures of the alloys show an increasing tendency for the basal poles to rotate away from the ‘normal direction’ towards the ‘rolling direction’. This can be attributed to the increased activity of the non-basal <c+a> slip mode, which can accommodate c-axis compression within individual grains, causing improved compressive ductility compared to pure magnesium. Betsofen et al. ⁶⁸ also found that the lithium alloying causes the formation of a prismatic rolling texture in the alloys as a result of the phase transformation of the lithium-based bcc phase into the magnesium-based hexagonal close-packed phase, which obeys the Burgers orientation relationships.

Al-Samman compared the deformation behaviour of Mg–4Li and AZ31 alloys. ⁶⁹ With respect to the room temperature formability examined in cold rolling and channel-die compression experiments, Mg–4Li alloy demonstrates a remarkable ductility enhancement over AZ31 alloy. During cold rolling, a total thickness reduction of 86% is easily achieved with no signs of cracks or mechanical instabilities. Correspondingly, AZ31 alloy exhibited a typical brittle response and failed at a total thickness reduction of 26% despite low reduction per each rolling pass (0·1 mm/pass), as shown in Fig. 19. The addition of Li seems to result in a congenerous deformation behaviour and texture evolution irrespective of the starting texture, which is completely different from that observed in AZ31, where the stress–strain mechanical response (flow curves) and texture evolution are strongly dependent on the starting texture. After cold rolling, Mg–4Li shows a basal texture with central peak intensity splitting, which is rotated towards the rolling direction. At room temperature, the addition of Li into magnesium seems to enhance the activation of <c+a>-pyramidal slip. With an increase of deformation temperature, the orientations seem to spread along TD forming an RD fibre texture component with maximum intensities at ±55° from RD. At a temperature of 400°C, the basal poles are completely rotated towards TD, developing a strong -prismatic texture. During deformation at 200°C, DRX takes place in Mg–4Li specimens mainly at grain boundaries and triple junctions. At 400°C, the deformed specimen reveals a coarse granular microstructure composed of grains with an average intercept length of ∼150 μm. Microstructure coarsening at 400°C is a consequence of grain growth. By contrast, AZ31 deformed under the same conditions does not reveal any signs of grain growth following recrystallisation.

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19.

Cold-rolling behaviour comparison between Mg–4Li and AZ31 alloys

Liu et al. ⁷⁰ investigated the microstructure and texture of Mg–3·3Li after ECAP with two different routes, A and Bc, respectively. Results show that two quite different and strong textures are formed in the alloy after two ECAP processes with the routes A and Bc. Recrystallisation occurs during both the ECAP processes, but the extent of recrystallisation is different for route A and Bc. ECAP can not only refine the grain size but also modify the texture type of the alloy. The combined effect of grain refinement and texture modification caused by ECAP can improve the strength and ductility of the alloy simultaneously.

Kumar et al. ⁷¹ studied the effect of hot-rolling process on the texture evolution of Mg–9%Li–5%Al–3%Sn–1%Zn (LATZ9531). The rolling textures for as-cast and as-rolled LATZ9531 are shown in Fig. 20 in terms of basal (0001) pole figures. It was found that strong basal texture is absent in the as-cast alloy (Fig. 20a ), and significant texture evolves after hot rolling at ∼573 K (Fig. 20b ). Random texture can be observed in the case of the as-cast alloy, as shown in Fig. 20a . On the other hand, the tilt of the basal-pole texture from the normal direction towards the transverse direction is observed in the as-rolled alloy. This strongly suggests that the sample is less prone to develop a basal texture, and the textures show even much weaker peak intensity, which is beneficial for further deformation because of availability of non-basal slip planes. The splitting of the basal poles about TD is attributed to an increased activity of the non-basal slip mode during compressive deformation. Central basal pole splitting along the sheet rolling direction is commonly observed during hot deformation of wrought commercial magnesium alloys such as AZ31; however, the splitting along the transverse direction is evident in the basal texture of the hot-rolled LATZ9531, which is attributed to the higher lithium and aluminium content, lowering c/a ratio.

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20.

The texture in terms of (0001) pole figures for as-cast and as-rolled LATZ9531 alloys: a as-cast; b as-rolled

Wu et al. ⁷² investigated the effects of the combined addition of Y and Nd on the room temperature compression texture of Mg–5Li–3Al–2Zn (LAZ532). The ODF sections for the alloys of LAZ532 and LAZ532–1·2Y–0·8Nd under different compression strains show that, when the strain is larger than 0·04, basal texture exists in LAZ532, while in LAZ532–1·2Y–0·8Nd, prismatic texture exists. The EBSD (Electron Backscattered Diffraction) analysis for LAZ532–1·2Y–0·8Nd at a strain of 0·04, as shown in Fig. 21, shows that some tension twins exist in the alloy (Fig. 21a ). The pole figure and inverse pole figure for the specimen show that the prismatic slip is activated. Because of the activation of prismatic slip and the precipitations of Al₂Y and Al₁₁Nd₃, the strength and plasticity of the alloy were improved as shown in Fig. 22.

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21.

EBSD analysis of Mg–5Li–3Al–2Zn–1·2Y–0·8Nd compression in room temperature: a orientation image map; b pole figure of (0001); c inverse pole figure

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22.

Compression stress–strain curves of the two alloys

Al₂Y_p reinforced Mg–Li composite

Compared with ceramic materials, some intermetallics have high specific strength, specific stiffness and high modulus, much superior plasticity, and metallic bond characteristics. Accordingly, these intermetallics can be utilised as one of the most promising reinforcements to fabricate composites with lightweight, high tensile and good integration mechanical properties. Al₂Y compound has low density (3·93 g cm⁻³). When compared to β–Mg–Li alloy, Al₂Y compound presents interesting properties: high melting temperature, high Young’s modulus, ⁷⁴ high hardness, and low coefficient of thermal expansion. Therefore, Al₂Y can serve as an excellent reinforcement in Mg–Li base alloy.

When 5 wt-% Al₂Y particulates (with an average size of less than 37·5 μm) are added into Mg–14Li–1Al (LA141) as a reinforcement (using stir casting with 700 rev min⁻¹ for 30 min), they distribute homogeneously in the alloy matrix. There is a clean interface between Al₂Y particles and the Mg–Li matrix, and no obvious elements diffusing to the interface. Accordingly, it is concluded that no particle/matrix interfacial reaction products exist in the Al₂Y_p/Mg–Li composite during the manufacturing process. The existence of the Al₂Y_p reinforcement significantly improves the mechanical properties of the alloy. The tensile strength, elastic modulus and hardness of the composite are higher than those of LA141 alloy by 45·3, 44·7 and 58·2%, respectively, while a good ductility property with an elongation of 7% can be maintained. ⁷⁵

When 15 wt-% Al₂Y particulates with a size range of 5–30 μm are added into Mg–12Li as a reinforcement in a vacuum non-consumable arc-melting furnace under an argon atmosphere, the Al₂Y particulates distribute homogeneously in the alloy matrix. A clean interface is formed between the Al₂Y particulates and the Mg–Li matrix. The hardness, shear strength and compression strength of the composite are higher than those of the alloy matrix by 150, 75 and 186%, respectively. ⁷⁶

δ-alumina fibre reinforced Mg–Li composite

Trojanova et al. ⁷⁷ prepared a series of Mg–Li based composites with the gas pressure infiltration into a fibrous perform (short alumina fibre). Matrix alloys of composites are Mg–xLi (x = 4,8,12), exhibiting the phase structure of hcp(α), bcc (β) and mixture (α+β) structures. The flow stresses of the composites are substantially higher than those of unreinforced alloy. The reinforcing effect of fibres and particles decreases with increasing temperature. The load transfer at the fibre/matrix interface and enhanced dislocation density are considered to be the most important strengthening mechanisms. Residual thermal stresses also play a significant role. There is a high probability that cross-slip and subsequent annihilation of dislocations cause softening. Local climb of dislocations in the vicinity of fibres supported by interface diffusion may also contribute to softening. When 10 vol.-% of short alumina fibre is added into Mg–8Li–xAl (x = 3,5), the strain hardening of the composite is characteristic of deformation at room temperature, as well as at 100°C. Recovery processes occur at higher temperatures. ⁷⁸ Similar to that in the Mg–xLi matrix composite, the contributions for the strengthening effect of the composite is the load transfer on the fibre–matrix interface and an increase of dislocation density.

In the Mg–Li/δ-alumina fibre composite, the formation of the interfacial bond in alumina fibre reinforced Mg-based composites is strongly promoted with lithium alloying. There are strong driving forces for redox reactions at the fibre/melt interfaces during the melt infiltration process producing Li⁺ that reacts immediately with adjacent alumina to form a metastable spinel-like compound δ(Li) structurally coherent with δ-Al₂O₃. Elemental aluminium as a byproduct of the redox alumina decomposition is dissolved in the Mg–Li matrix causing its additional alloying as demonstrated schematically below: ^79,80 (1) (2) (3) where solid-state reaction (2) is essential for the formation of the interfacial bond.

Mg–Li alloy reinforced by short δ-alumina fibre shows an anisotropy of thermal expansion. Kudela et al. ⁸¹ prepared Mg–4Li and pure Mg matrix composites reinforced by short δ-Al₂O₃ fibres with melt infiltration at 730°C/30 s under argon pressure (6 MPa), referred to as Mg4Li/SF and Mg/SF. Prepared composites contained 10 vol.-% fibres of mean length of ∼300 μm and mean diameter of ∼3 μm. The fibres were arranged randomly in planes parallel to each other. This gives rise to two distinct fibre configurations with fibre planes either parallel (‘in-plane’) or perpendicular (‘cross-plane’) to the axis of the composite rod (6 mm in diameter, 50 mm in length). The measurements for the coefficiency of thermal expansion (CTE) showed that the in-plane Mg/SF expands less than the in-plane Mg4Li/SF. Expansion curves of both composites were located well below those of unreinforced Mg and Mg4Li matrixes, demonstrating a thermal expansion reduction because of the fibres. In cross-plane, the CTE of Mg/SF was larger than that of Mg4Li/SF, while in in-plane, the CTE of Mg/SF was lower than that of Mg4Li/SF. The temperature dependence of CTE of Mg4Li alloy and both in-plane and cross-plane Mg4Li/SF composites is shown in Fig. 23. When the temperature is lower than 300°C, both in-plane and cross-plane CTE curves of Mg4Li/SF are less than that of Mg4Li, and the cross-plane CTE curve is higher than the in-plane CTE curve of Mg4Li/SF. When the temperature is larger than 300°C, the cross-plane CTE curve of Mg4Li/SF is higher than that of Mg4Li, and the in-plane CTE curve of Mg4Li/SF decreases with temperature.

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23.

Temperature dependence of CTE of Mg4Li alloy and both in-plane and cross-plane Mg4Li/SF composites

SiC reinforced Mg–Li composite

The enforcement of SiC_w in Mg–Li binary alloys improves the damping capacity, and it is shown that the high Li content is beneficial to improvement of damping capacity, as shown in Fig. 24. ⁸² The high damping capacity of SiC_w/MgLiAl composites is attributed to the increased interfacial defect and the higher dislocation density induced by the addition of SiC_w. Compared with the Mg–Li matrix, higher tensile strength and higher modulus are obtained. Moreover, the specific strength and specific modulus are also increased. Research on the interface between SiC_w and Mg–Li matrix shows that a clear interface can be obtained. SiC_w connects with the matrix in (111) and forms 70·5° or 109·5° stages on the whisker surface in (111). ⁸³

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24.

Comparison of the damping capacity for MgLi4Al1, SiCw/MgLi4Al1 and SiCw/MgLi8Al1 (ϵ = 1×10⁻⁴, f = 1 Hz)

B, B₂O₃ and B₄C reinforced Mg–Li composite

In B particle reinforced Mg–Li base alloys, both a high strength and a high ductility can be obtained because of the existence of Li, which can counteract the negative effect of the existence of hard B particles on the ductility. ⁸⁴ A high specific stiffness can also be expected because of the low density of Li element. The existence of B reinforcement in the Mg–14Li matrix is contributed to the improvement of creep behaviour at 230–280°C. The creep strength of 30 vol.-%B/Mg–14Li is increased by a factor of eight over the Mg–14Li matrix. ⁸⁵

The addition of B₂O₃ in Mg–Li melt causes the following reaction: ⁸⁶ (4) This reaction can take place at 641·3–679·2°C to produce MgO and B reinforcements. To obtain good mechanical properties in the alloy, a relatively low temperature (less than 800°C) and a relatively rapid solidification are required to avoid the existence of coarse MgO.

Gonzalez-Doncel et al. ⁸⁷ used the foil metallurgy method, a thermo-mechanical process based on low temperature press bonding of cold-rolled and annealed foils, to prepare a fine-grained Mg–9Li laminates and a particulate composite of Mg–9Li containing 5 wt-% B₄C. As for Mg–Li base composite, this method is potentially more attractive than powder metallurgy processing because contamination is reduced as a result of the low surface-to-volume ratio in the foils. The total amount of reduction in foil thickness during cold rolling determines the final grain size of the composite alloy. The total reduction of 170∶1 corresponds to a fine grain size of 2 μm. The room temperature ductility of the composite is higher than that of the as-cast alloy by 50%, and the specific stiffness is 3·16×10⁶ m³, about 22% above that of commercial aluminium and titanium alloys. The composite possesses superplasticity at 150–200°C with a stress exponent of 2. The flow stress in compression is about two to three times higher than in tension. This can be attributed to the greater ease of GBS in tension than in compression. The activation energy for superplastic flow is 55 kJ mol⁻¹ and is essentially independent of the mode of testing. ⁸⁸ It is postulated that superplastic flow of the Mg–9 wt-%Li–5 wt-%B₄C particulate composite occurs by GBS controlled by grain boundary diffusion.

The pre-treatment of alkaline solution for B₄C particles can not only reduce the agglomerates and produce uniform distribution of the B₄C particles in the composite but can also enhance the wetting ability between B₄C particles and Mg–Li matrix. The tensile strength of B₄C/Mg–Li composite increases 21·53% compared with that without alkaline solution treatment, while the elongation decreases by only 2·29%. ⁸⁹

Wu et al. ⁹⁰ prepared B₄C_p/Mg–8Li–1Zn and B₄C_p/Mg–8Li–1Al–1Y composites using hot-extrusion solid-state composite processing. With the optimised parameters, the deformation effects and the migration of α phase between the foils were improved, and the amount and size of foil gaps decreased. The bonding force between foils was improved, and the oxidation of foils was lowered. The results of tensile test showed that the strengths of the B₄C_p/Mg–8Li–1Zn and B₄C_p/Mg–8Li–1Al–1Y composites were increased after hot-extrusion solid-state composite processing (238 and 257·23 MPa, respectively).

Steel fibre reinforced Mg–Li composite

Steel wool (bcc structure) can be used as reinforcement in Mg–Li matrix with the infiltration casting process. The use of discontinuous, chopped fibres, rather than a layered, continuous fibre preform, can somewhat improve the deformation characteristics of the composite. As for hcp/bcc Mg–4Li/27 vol.-%steel composite, it has limited formability because of the severity of internal cracking, shown in Fig. 25. ⁹¹ The strength of the Mg–4Li/27steel increases from 163 MPa as-cast to 332 MPa at a true strain of 3·9. Strengths are lower than predicted from ROM calculations. As for bcc/bcc Mg–12Li–2Nd/21vol.-% steel composite, it has a good formability (reaching a true strain of 9·2) when swaged at room temperature. The tensile strength increases from 180 to 326 MPa at a strain of 4·3. Strengths are also lower than predicted from ROM calculations because of the presence of fibres considerably larger than the average size measured stereologically. ⁹² The bcc/bcc material shows the potential for producing highly formable composite sheet materials by combining a ductile bcc matrix with a bcc reinforcing fibre. Figure 26 shows the different formability of Mg–4Li/27 vol.-%steel and Mg–12Li–2Nd/21vol.-%steel.

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25.

Internal cracking on hcp/bcc Mg–4Li/27 steel composite during deformation

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26.

Comparison of the rolling characteristics of hcp/bcc Mg–4Li/27steel (top) and bcc/bcc Mg–12Li–2Nd/21steel (bottom)

Al/Mg–Li cladding composite

It is well known that Mg–Li alloys possess poor corrosion resistance, while Al alloys have excellent corrosion resistance because of the formation of a thermodynamically stable Al₂O₃ protective surface oxide scale. In addition, various surface-treatment techniques are already established industrially for Al alloys. Therefore, the cladding of Mg–Li and Al alloy plates is considered to be promising to improve the corrosion resistance of Mg–Li alloy plates.

Matsumoto et al. ⁹³ fabricated pure Al/Mg–Li/pure Al three-layered plate by cold clad rolling and annealing. When annealing at or above 523 K, the reaction between layers occurs. At the pure Al side, the reaction phase is Li-dissolved Al₃Mg₂. At the side of Mg–Li side, it is highly concentrated Mg phase. In 180° bending test, fracture of the clad plate annealed at or above 523 K is initiated at the reaction phases at the joint interface. The clad plates annealed at 423 and 473 K have no reaction phase at the joint interface and exhibit excellent bend formability. In tensile testing, the clad plates annealed at or above 423 K show good ductility over 20%. During annealing at or below 473 K, interface debonding is not observed after fracture. During annealing at or above 523 K, multiple cracking is initiated in Al₃Mg₂ and subsequently a main crack always propagates throughout fragmented Al₃Mg₂, not along the interfaces, thereby leading to macroscopic debonding. When Al–Mg/Mg–Li/Al–Mg composite sheet was fabricated by clad rolling at room temperature, and the composite sheet was annealed at various annealing conditions after clad rolling, at Al–Mg alloy side, the reaction phases were Al₃Mg₂ and LiMgAl₂. ⁹⁴ At Mg–Li side, the reaction phase is still a highly concentrated Mg phase. The width of the reaction phases is broadened with annealing by lattice diffusion control, where the growth constant K (m²/s) of LiMgAl₂, Al₃Mg₂ and high Mg content phase at 573 K was calculated to be 1·3×10⁻¹⁵ m² s⁻¹, 9·0×10⁻¹⁵ m² s⁻¹ and 3·9×10⁻¹⁵ m² s⁻¹, respectively. Crack propagation is observed only through the reaction phases, leading to interface debonding in 90° bending test and tensile test. It is concluded that the three-layered clad sheet shows good ductility and high strength on annealing at low temperatures such as 423 and 473 K, where the formation of reaction phases is suppressed.

Zu et al. ⁹⁵ studied the preparation processes for Mg–9Li–1Zn/Al clad composite. The critical reduction of Mg–9Li–1Zn alloy/Al cladding plate is about 50%, and the best reduction is about 60–65%. After annealing, α phase can be spheroidised, and the complete recrystallisation occurs after annealing at 300°C for 3 h. At the same time, both the microstructure and properties of the composite plate are optimised.

Codeposition of Mg–Li–Al alloys

Mg–Li–Al alloys have become very attractive in recent years because of their low densities and superior mechanical properties. ¹¹⁶ Yan et al. have investigated the codeposition of Mg–Li–Al alloys from LiCl–KCl–MgCl₂–AlCl₃ melts by employing a series of electrochemical techniques. The Mg–Li–Al alloys with different phases were prepared by potentiostatic and galvanostatic electrolysis. Figure 28 shows chronopotentiograms measured on a molybdenum electrode in LiCl–KCl–MgCl₂ (0·525 mol kg⁻¹) melts containing 0·075 mol kg⁻¹ AlCl₃ at different current intensities. ¹¹⁷ At a cathodic current lower than −20 mA (current density −0·052 A cm⁻²), the curves exhibit two potential plateaus (plateaus 1 and 2), which are associated with the reduction of aluminium (plateau 1) and magnesium (plateau 2) ions to metals. When the current reaches −180 mA (−0·47 A cm⁻²), a third plateau appears caused by the reduction of lithium ions. At this current intensity, codeposition of Mg, Li and Al occurs. It is obvious that the potential ranges for the deposition of Mg, Li and Al are the same as those observed in the cyclic voltammograms (CVs). Ye et al. ¹¹⁸ also attempted to codeposit Mg and Li ions at an Al electrode to form Al–Li–Mg alloys in LiCl–KCl melts containing different concentrations of MgCl₂ at 893 K. They prepared Al–Li–Mg alloys by galvanostatic and potentiostatic electrolysis and claimed that the lithium and magnesium contents can be controlled by MgCl₂ concentrations by electrolytic parameters.

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28.

Chronopotentiograms obtained at different current intensities on a molybdenum electrode (0·3847 cm²) in the LiCl–KCl–MgCl₂ (5 wt-%) melts containing 1 wt-% AlCl₃ at 670°C

The ICP analyses of all samples obtained by potentiostatic and galvanostatic electrolysis are listed in Tables 5 and 6, respectively. During potentiostatic electrolysis, the more negative the applied potential in the LiCl–KCl–MgCl₂ (0·525 mol kg⁻¹) melts with equivalent AlCl₃ concentration, the higher the lithium content of Mg–Li–Al alloy. In addition, the aluminium content of Mg–Li–Al alloys increases with increasing AlCl₃ concentrations in LiCl–KCl–MgCl₂ melts. Under galvanostatic electrolysis, the lower MgCl₂ concentration in the LiCl–KCl melts with equivalent AlCl₃ concentration at a constant current intensity, the higher the lithium content of the Mg–Li–Al alloys. The aluminium content of Mg–Li–Al alloys increases with an increase in AlCl₃ concentrations in LiCl–KCl–MgCl₂ melts. These results indicate that the lithium and aluminium contents of Mg–Li–Al alloys are adjustable simply by changing the concentrations of MgCl₂ and AlCl₃ and the electrolytic parameters.

Codeposition of Mg–Li–Zn alloys

Zn is an important alloying element in Mg base alloys. Yan et al. ¹¹⁹ studied the mechanism of codeposition of Mg–Li–Zn alloys. Figure 29 shows typical CVs obtained at a molybdenum electrode before and after the addition of 8 wt-% MgCl₂ and 1 wt-% ZnCl₂ in LiCl–KCl melts at 943 K (The dotted curve represents the voltammogram before the addition of MgCl₂ and ZnCl₂.). Only one couple of cathodic/anodic (C/C′) signals were observed, which corresponds to the deposition and dissolution of liquid Li. The solid curves show the voltammograms measured at different cathodic limits after the addition of MgCl₂ and ZnCl₂. Peaks A and B in the negative-going scan are ascribed to the formation of a three-dimensional (3D) phase of Zn and Mg, respectively, because the deposition potential of Zn ions is more positive than that of Mg ions in a chloride system. Peaks B′ and A′ in the positive-going scan correspond to the dissolution of Mg and Zn deposit, respectively. Before the addition of MgCl₂ and ZnCl₂, reduction of Li (I) starts at about −2·25 V on a molybdenum electrode; after the addition of MgCl₂ and ZnCl₂, the reductions of ions of Zn, Mg, and Li start at approximately −0·70, −1·55, and −2·00 V, respectively. In this way, the potential of Li metal deposition after the addition MgCl₂ and ZnCl₂ is more positive than the Li metal deposition before the addition. The potential shift is because of a lowering of activity of the deposited metal (Li) in a foreign substrate. The foreign substrate is probably Mg–Zn alloy pre-deposited on the molybdenum electrode. Galvanostatic electrolysis was carried out in LiCl–KCl melts containing MgCl₂ and ZnCl₂ with different concentrations on molybdenum electrodes at 943 K. Figure 30 shows the complicated XRD patterns of Mg–Li–Zn alloy samples obtained by galvanostatic electrolysis from the LiCl–KCl melts containing 5–10 wt-% MgCl₂ and 1–3 wt-% ZnCl₂ at 6·21 A cm⁻² for 2 h. As seen from the XRD patterns, the Mg–Li–Zn alloys are composed of α+Mg₇Zn₃, (α+β)+Mg₇Zn₃, and β+LiZn phases. In patterns a–c, all Mg–Li–Zn alloys contain the Mg₇Zn₃ phase, which proves that the first formed Mg–Zn alloy is Mg₇Zn₃. However, a new LiZn phase occurs in pattern d when the concentration of MgCl₂ is 5 wt-% in LiCl–KCl–ZnCl₂ (1 wt-%) melts. Moreover, the lithium content of Mg–Li–Zn alloys increases (from α, α+β to β phase) with a decrease of MgCl₂ concentrations in the LiCl–KCl–ZnCl₂ (1 wt-%) melts at constant current intensity.

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29.

Typical cyclic voltammograms (CVs) of the LiCl–KCl melts before (dotted line) and after (solid line) the addition of 8 wt-% MgCl₂ and 1 wt-% ZnCl₂ on molybdenum electrodes at 670°C

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30.

X-ray diffraction (XRD) patterns of deposits obtained by galvanostatic electrolysis on Mo electrodes (S  =  0·322 cm²) in the LiCl–KCl melts with a 10 wt-% MgCl₂ and 1 wt-% ZnCl_2; b 9 wt-% MgCl₂ and 1 wt-% ZnCl₂; c 9 wt-% MgCl₂ and 3 wt-% ZnCl₂; d 5 wt-% MgCl₂ and 1 wt-% ZnCl₂ at 2 A for 2 h

Codeposition of Mg–Li–Ca alloys

The addition of low-cost Ca into magnesium-based alloys can raise the oxidising combustion temperature of magnesium-based alloy, refine the microstructure, and improve the mechanical properties at room temperature and heat resistance at elevated temperature. ¹²⁰ Yan et al. have successfully prepared Mg–Li–Ca alloys by codeposition of Mg, Li and Ca from LiCl–KCl–MgCl₂–CaCl₂ melts. ¹²¹ According to Faraday’s law, the mass of the deposited substance is related to I and t. If the authors want to obtain a larger mass Mg–Li–Ca alloy in a relatively short time, they have to perform the electrolysis at a more negative potential or current density. Under potentiostatic electrolysis, however, the current often exceeded the current limit of the electrochemical workstation (−2–2 A, −4–4 V) with the growth of the electrode surface (growth of the nucleation of the M g–Li–Ca alloy on a Mo electrode). Therefore, the experiment cannot be continuously completed because of an interruption in the measurement. Certainly, this problem can be overcome when an updated instrument with a wider current range is employed. This is because there was a limit to the increase of the electrode surface. The nucleation of Mg–Li–Ca alloy on Mo electrode grew to some extent and then the liquid Mg–Li–Ca alloy departed from the Mo electrode because of the effect of surface tension. Under galvanostatic electrolysis, this problem did not occur because of a wide potential range of the electrochemical workstation. Therefore, galvanostatic electrolysis was carried out using a much higher current density than the onset one for codeposition of Mg, Li and Ca in LiCl–KCl–CaCl₂–MgCl₂. The microstructures of the Mg–Li–Ca alloys by codeposition from LiCl–KCl–CaCl₂ (1 wt-%) melts containing 9 and 8 wt-% MgCl₂ exhibit typical α and β phases, respectively, as given in Fig. 31a and b . The average grain sizes of the α and β Mg–Li–Ca alloys are about 100 and 70 μm, respectively. In order to determine the distribution of the elements Mg and Ca in the Mg–Li–Ca alloy, mapping analysis of the elements was employed. Figure 32 shows a scanning electron microscopy (SEM) and EDS mapping analyses of the Mg–Li–Ca alloy by codeposition from LiCl–KCl–CaCl₂ (1 wt-%) melts containing 8 wt-% MgCl₂. The element Mg distributes homogeneously throughout the Mg–Li–Ca alloy. However, Ca distribution is not uniform and mainly disperses along the grain boundaries.

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31.

Optical micrographs of the Mg–Li–Ca alloys by codeposition from LiCl–KCl–CaCl₂ (1 wt-%) melts containing a 9 wt-% and b 8 wt-%MgCl₂

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32.

Scanning electron microscopy (SEM) and EDS mapping analyses of the Mg–Li–Ca alloy by codeposition from LiCl–KCl–CaCl₂ (1 wt-%) melts containing 8 wt-% MgCl₂

Codeposition of Mg–Li–Mn alloys ¹²²

Galvanostatic electrolysis was carried out in LiCl–KCl melts containing MgCl₂ and MnCl₂ with different concentrations on molybdenum electrodes at 893 K. Figure 33 shows the XRD patterns of Mg–Li–Mn alloy samples obtained by galvanostatic electrolysis from the LiCl–KCl melts containing 6–11 wt-% MgCl₂ and 2 wt-% MnCl₂ at 7·14 A cm⁻² for 2 h. The observed peaks were identified as α-Mg, β-Li, and Mn phase. As seen from the XRD patterns, the Mg–Li–Mn alloys are composed of α-Mn, (α+β)+Mn, and β+ Mn phases. In patterns a–c, all Mg–Li–Mn alloys contain the Mn phase, which proves that intermetallic compound was not formed in Mg–Li–Mn alloys at this temperature. According to the phase diagram of the Mg–Mn and Li–Mn system, the concentration of Mn will be small in Mg–Li alloys since the solubility of Mn in Mg is around 0·8 at-% and Mn is insoluble in Li metal at 893 K. Moreover, the lithium content of Mg–Li–Mn alloys increases (from α, α+β to β phase) with a decrease in MgCl₂ concentrations in the LiCl–KCl–MnCl₂ (2 wt-%) melts at constant current intensity. Therefore, under the condition of electrolysis, the lithium content and phase composition of Mg–Li–Mn alloys are adjustable simply by changing concentrations of MgCl₂ and the electrolytic parameters.

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33.

X-ray diffraction (XRD) patterns of deposits obtained by galvanostatic electrolysis on Mo electrodes (S = 0·28 cm²) in the LiCl–KCl melts with a 11 wt-% MgCl₂ and 2 wt-% MnCl₂; b 8 wt-% MgCl₂ and 2 wt-% MnCl₂; c 6 wt-% MgCl₂ and 2 wt-% MnCl₂ at 2·00 A for 2 h at 893 K

Codeposition of Mg–Li–Sb alloys

Sb is one of alloying elements of Mg alloys. Adding a small amount of Sb could improve mechanical properties, thermal stability and grain refinement in comparison with the base alloy. ¹²³ Wei et al. ¹²⁴ investigated the preparation of Mg–Li–Sb alloys by galvanostatic electrolysis in LiCl–KCl melts containing MgCl₂ and SbCl₃ with different concentrations on a molybdenum electrode at 673 K for 2 h. Owing to the low electrolysis temperature, after the process of electrolysis, the temperature of the salts was increased to 973 K after holding for 0·5 h. Li content of Mg–Li–Sb alloys increases with a decrease in MgCl₂ concentration in LiCl–KCl melts. The higher the SbCl₃ concentration in LiCl–KCl melts, the higher the Sb content of Mg–Li–Sb alloy can be obtained. According to these results, Sb and Li contents of Mg–Li–Sb alloys can be adjusted by changing MgCl₂ and SbCl₃ concentrations in LiCl–KCl melts. Figure 34 shows the XRD patterns of alloys listed in Table 7. With an increase of Li and Sb content, the amount of Li₃Sb increased. When Li content is 12·90 wt-% and Sb content is 2·70 wt-%, Mg₃Sb₂ phases appear in the alloy.

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34.

X-ray diffraction (XRD) patterns of deposits obtained by galvanostatic electrolysis on a Mo electrode (S = 0·322 cm²) in the LiCl–KCl melts containing different concentrations of MgCl₂ and SbCl₃ at 673 K for 2 h, and then enhanced melts temperature to 973 K holding 0·5 h

Codeposition of Mg–Li–REs (rare earths) alloys

Most commercial Mg–Li–RE alloys are prepared by directly mixing and melting pure magnesium, lithium and rare earth metals. This conventional method results in high energy consumption for industrial production of pure magnesium, lithium and rare earth metals. Specifically, pure RE metals are extremely expensive because it is too difficult to extract and purify them. In recent years, electrochemical codeposition has been used to prepare RE alloys. A series of Mg–Li–RE alloys have been successfully prepared.

Variable valence REs (Sm, Eu and Yb) cannot be directly electrolysed from LiCl–KCl melts since the reduction potential of RE(II) to RE(0) is even more negative than that of Li(I) to Li(0). It is well known that the depolarisation effect of REs with Mg is larger than that of Li. Therefore, codeposition of REs and Mg ions is an effective way to obtain REs by electrolysis. Ye et al. ^125,126 have investigated electrodeposition of Mg–Yb alloy film and Mg–Li–Yb alloys at solid magnesium cathode in the molten LiCl–KCl–YbCl₃ system at 773 K.

Figure 35 shows typical CVs obtained at a Mo electrode (curve 1) and an Mg electrode (curves 2 and 3) before and after the addition of 2 wt-% YbCl₃ in LiCl–KCl melts at 743 K. The dotted curve 1 represents the voltammogram before the addition of YbCl₃ at a Mo electrode. A sharp increase in cathodic current from approximately −2·33 V (versus Ag/AgCl) was observed. The cathodic signal B can be ascribed to the deposition of Li. The dash curve 2 shows the voltammogram before the addition of YbCl₃ at an Mg electrode. A cathodic current was seen from about −2·12 V. Since this potential value is more positive than that of Li metal deposition, the cathodic current is thought to be caused by the formation of Mg–Li alloys. The potential shift is because of a lowering of activity of the deposited metal (Li) in the Mg phase. After adding 2 wt-% YbCl₃, peak A ₁ associated with the reduction of Yb(III) to Yb(II) ions was first detected at about −0·77 V in solid curve 3. Afterwards, a cathodic current was observed. The cathodic current A ₂ probably corresponds to the formation of Mg–Yb alloys. Finally, the cathodic signal B observed in the cathodic limit can be ascribed to the formation of Mg–Li–Yb alloys. In the anodic direction, peaks B′ and A ₂′ correspond to the dissolution of Li (0) and Yb (0) from the Mg–Li–Yb alloys, respectively. Afterwards, anodic peak A ₁′ can be ascribed to the oxidation of Yb(II) to Yb(III) ions. Potentiostatic electrolysis experiments were conducted in the LiCl–KCl–YbCl₃ (2 wt-%) melts for 3 h using Mg electrodes as a working electrode at −1·80 V (sample 1) and −2·40 V (sample 2). Figure 36 shows the XRD patterns of samples 1 and 2. The diffraction peaks for sample 1 could be assigned to α-Mg phase, indicating the Yb and Li contents in Mg–Li–Yb alloys were relatively low. However, typical α+β phase Mg–Li–Yb alloys were formed when the applied potential is at −2·40 V. The observed peaks for sample 2 are identified as (α+β)+Mg₂Yb phases.

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35.

Cyclic voltammograms (CVs) of LiCl–KCl melts before and after the addition of 2 wt-% YbCl₃ at a Mo electrode (S = 0·104 cm²; curve 1) and an Mg electrode (S = 0·457 cm²; curves 2 and 3) at 743 K. Scan rate: 0·1 V s⁻¹

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36.

X-ray diffraction (XRD) patterns of samples 1 and 2 (followed by a post-thermal treatment at 973 K) at −1·80 and −2·40 V (versus Ag+/Ag) obtained at Mg electrodes (S = 0·457 cm²) in the LiCl–KCl–YbCl₃ (2 wt-%) melts by potentiostatic electrolysis for 3 h at 743 K, respectively

Zhang et al. ^127–129 have investigated the preparation of Mg–Li–RE (La, Gd and Yb) alloys from RE₂O₃ in the LiCl–KCl(–KF)–MgCl₂ melts. Galvanostatic electrolysis was carried out in a molten KCl–LiCl–MgCl₂–RE₂O₃–KF system on molybdenum electrodes at 923–973 K. In the beginning stage of electrolysis, compared with the higher concentration of LiCl in the melts, MgCl₂ reached limited diffusion current density and deposited first. The temperature of 923–973 K was higher than the melting point of metal magnesium (923 K). So, the liquid magnesium cathode was around the molybdenum wire. And then Li⁺ and RE³⁺ began to deposit on the surface of the pre-coated liquid magnesium cathode.

From Table 8, the lower the MgCl₂ concentration in the LiCl–KCl–KF melts, with equivalent La₂O₃ concentration at a constant current intensity, the higher the lithium content of the Mg–Li–La alloys. In the case of neglecting La deposition (the concentration of La₂O₃ in the melts is very low), the molar ratio of deposited Mg and Li is proportional to the ratio of their current densities. The limited diffusion current density of MgCl₂ is proportional to its concentration in the melts. Therefore, Mg content in the alloy decreases with decreasing MgCl₂ concentration in the melts. The authors can also find that the lanthanum content of Mg–Li–La alloys increase with increasing the La₂O₃ concentrations in KCl–LiCl–MgCl₂–La₂O₃–KF melts.

In order to examine the uniformity of the elements Mg and Yb distributed in Mg–Li–Yb alloys, elemental mapping analysis was employed. Figure 37 shows SEM morphology (Fig. 37a ) and EPMA area analysis (Fig. 37b and c ) of Mg–Li–Yb alloys prepared by galvanostatic electrolysis with 3 A, 60 min at 660°C in the melts (LiCl∶KCl∶KF∶MgCl₂ Yb₂O₃ = 40∶40∶10∶9·5∶0·5, wt-%) on Mo electrodes. The grey zone in the SEM corresponds to α-Mg phase. Figure 37b shows that the element Mg distributes homogeneously in the Mg–Li–Yb alloys. The distribution of Yb is not uniform in the alloy and is mainly distributed in the grain boundaries of α-Mg, as presented in Fig. 37c .

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37.

Scanning electron microscopy (SEM) and EPMA mapping analysis of the Mg–11·8Li–8·4Yb alloy obtained by galvanostatic electrolysis with 3 A, 60 min at 660°C in the melts (LiCl∶KCl∶KF∶MgCl₂∶Yb₂O₃ = 40∶40∶10∶9·5∶0·5, wt-%) on Mo electrodes a SEM; b and c Area analysis of Mg and Yb in Mg–Li–Yb alloy

Zhang et al. ^130–133 have also prepared Mg–Li–RE (Y, Sm and Er) alloys from RECl₃/K₃RECl₆ in the LiCl–KCl–MgCl₂ melts. Figure 38 shows the CVs of the LiCl–KCl melts before and after the addition of MgCl₂ and SmCl₃ on molybdenum electrodes at 903 K. The dotted line corresponds to a CV of the LiCl–KCl melts in the absence of MgCl₂ and SmCl₃. A sharp increase in cathodic current from –2·40 V is observed. The cathodic signal C can be ascribed to the deposition of Li because no alloys or intermetallic compounds exist for the Mo–Li binary system at 903 K. In the reverse scan, an anodic peak C′ corresponding to the dissolution of Li is observed. The solid curves show the voltammograms measured after the addition of MgCl₂ and SmCl₃. Two cathodic/anodic current peaks observed at approximately –0·90 and –0·75 V are attributed to the reduction of Sm(III) and the oxidation of Sm(II), respectively. In the negative-going scan, peak B is observed from approximately –1·80 V, which corresponds to the reduction of Mg(II). The deposition of Li(I) is observed at a more positive potential after the addition MgCl₂ and SmCl₃ than that before the addition. The potential shift is caused by the formation of Mg–Li alloys. The redox couple Sm(II)/Sm(0) is not observed in this experiment because the reduction of Sm(II) is at a high negative value close to that of the lithium reduction. It is probably hidden by a larger scaled cathodic current of Li deposition.

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38.

The cyclic voltammograms (CVs) of the LiCl–KCl melts before (dotted line) and after (solid line) the addition of MgCl₂ and SmCl₃ (CMg(II) = 1·73×10⁻⁴ mol cm⁻³, CSm(III) = 1·359×10⁻⁴ mol cm⁻³) on molybdenum electrodes at 903 K (903 K). Scan rate: 0·1 V s⁻¹

From Table 9, under galvanostatic electrolysis, the samarium content of Mg–Li–Sm alloys increases with increasing SmCl₃ concentrations in LiCl–KCl–MgCl₂ melts. With equivalent SmCl₃ concentration at constant current intensity, the longer the electrolytic time, the higher the lithium content in Mg–Li–Sm alloys. Based on these results, it can be concluded that the lithium and samarium contents in Mg–Li–Sm alloys can be adjusted by changing the electrolytic time and concentration of SmCl₃.

Chromate conversion coatings

Chromate conversion coating is a relatively common conversion coating on Mg–Li alloys, which can be used as pre-treatments before a final sealing process or as ‘post-treatments’ after a plating process to improve corrosion resistance, paint or to provide a decorative finish or adhesive bonding properties. One of the main disadvantages of chromate conversion coatings is the toxicity of the treatment solutions. Accordingly, the chromate solution have been restricted and forbidden in industries in many countries.

When applying a conversion coating, there are a few general rules that should be followed. First, the Mg–Li alloy substrates having fine-grained microstructure respond to chromating. Second, a proper pre-treatment of the surface is necessary to achieve the best coatings. Third, after chromating, the Mg–Li alloys substrates should be properly rinsed in order to prevent the residual acid or base from reacting with the coating. Finally, the coatings should be air dried at a temperature of 70°C for a few minutes.

The deposition rate of the conversion coatings in pure chromic acid solutions is very slow. ¹⁵³ Some anions can act as a catalyst for deposition in solution, ¹⁵⁴ including formate, acetate, chloride, sulphate, nitrate, fluoride, sulphamate and phosphate ions. ¹⁵⁵ The pH of the solution is the most important factor in controlling the formation of chromate coatings on Mg–Li alloys. ¹⁵⁶ The chromium coatings provide corrosion protection by presenting a non-reactive barrier to the environment. ^155–157 The protection afforded by the chromium coating is proportional to the coating thickness. In order to maintain its protective properties, the coating should be subjected to low temperature. This would increase the coating thickness and achieve good self-healing ability of the coating. At lower temperature, the coating retains its self-healing characteristics as long as it remains in its hydrated form. The stability of the coatings at higher temperature can be improved by sealing or painting on top of the conversion layer. The hexavalent chromium is reduced during corrosion to form trivalent chromium in order to terminate the oxidative attack.

A study on chromate conversion coating on Mg–Li alloys ¹⁵⁷ suggested that the coatings around 8–11 μm in thickness present excellent adhesion even under humidity and thermal cycling tests. The optical and paint-base properties of the coating were also unaffected by humidity and thermal cycling tests.

Phytic acid conversion coatings

Phytic acid conversion treatment is a novel protection technique for Mg–Li alloys, owing to its non-toxicity and peculiar structure. ^158–160 The phytic acid conversion coating can slow down the corrosion of Mg–Li alloys. Many scholars have studied the growth and formation of a conversion coating on Mg–Li alloys in a phytic solution. The coating reduced the corrosion current density, the corrosion potential and the hydrogen evolution rates.

The active groups of phytic acid can react with metal ions, such as Al³⁺ and Mg²⁺, and form stable chelate compounds on the surface of Mg–Li alloys. Because of phytic acid containing six phosphate carboxyl groups and hydroxyl groups, the phytic acid conversion coating may react with organic coatings and increase the adherence between the organic coatings and the Mg–Li alloys. Phytic acid conversion coatings were uniform and with white, flower-like deposit formed on Mg–Li alloy surfaces. ¹⁶¹ The coating improved the corrosion resistance of Mg–Li alloys. The optimum process parameters were confirmed as follows: pH value of the solution is 6, concentration of phytic acid is 20 g/L, treating time is 10 min and treating temperature is 35°C. The pH value of the solution was the main factor influencing the growing of the phytic acid conversion coating. The conversion coating consisted of Mg, Al, O, P and C. Chelate compounds were formed between phytic acid and Mg²⁺ or Al³⁺.

Phosphate–permanganate conversion coatings

Highly stable chromate conversion coatings on Mg–Li alloys have been obtained, ¹⁶² but chromate is not friendly to the environment. Consequently, there is a need to develop new environment-friendly conversion treatments for Mg–Li alloys. Many researchers have studied the growth of a chrome-free conversion coating on Mg–Li alloys in a solution of phosphate–permanganate. The phosphate–permanganate conversion coatings have been shown to have corrosion resistance comparable to chromate treatments. The research on phosphate–permanganate treatment of Mg–Li alloys ¹⁶³ using a solution containing sodium phosphate and potassium permanganate has shown that a non-powdery, homogeneous and uniform coating can be achieved. The pH and phosphate concentration of the conversion solution were the most effective on the quality of the final coating.

A selected solution of phosphate and permanganate resulted in the formation of an adherent and continuous layer on the Mg–Li alloys substrates. The corrosion resistance, the composition and the morphology of the phosphate–permanganate coating were examined. The phosphate–permanganate conversion coating on the Mg–Li alloys was uniform and environmentally friendly, with a phosphate–permanganate solution. ¹⁶⁴ The conversion coating was obtained at the following operating conditions: KMnO₄ 40 g/L, KH₂PO₄ 50 g/L, 328 K, and 20 min. The results of immersion and electrochemical tests confirmed that the phosphate–permanganate conversion coating improved the corrosion resistance of Mg–Li alloys. XPS and EDX analyses confirmed that the phosphate–permanganate conversion coating were mainly composed of elements Mg, O, K, P and Mn, and the corrosion behaviour was the same for both chromate and phosphate–permanganate coatings.

After the pre-treatment in a solution of phosphate–permanganate, Mg–Li alloy samples exhibit good paint adhesion. The best quality conversion coating was found to be the control of the pH, which is the most important factor in pre-treatment. The permanganate can act as a catalyst without depositing metallic manganese on the Mg–Li alloys surface. The phosphate–permanganate coatings were shown to have good paint base performance and corrosion resistance. The phosphate–permanganate coatings were composed of an agglomerate of well-formed crystalline phosphate compounds.

Rare earth conversion coatings

So far, investigations on rare earth conversion coating treatments for the very active Mg–Li alloys have been few. Recently, rare earth conversion treatments were studied in order to find effective chromate-free treatments for Mg–Li alloys. ^165–168 A lanthanum-based conversion coating for Mg–Li alloys by immersion in 5 g/L La(NO₃)₃ solution. The solution used is environmentally friendly. ¹⁶⁹ A rare earth conversion coating was studied with cracked and uniform morphology, which was formed on the Mg–Li alloys from a solution of lanthanum nitrate and cerium nitrate. The surface morphologies before and after corrosion in 3·5 wt-% NaCl solution are shown in Fig. 44. A homogeneous and cracked layer was observed on the surface of Mg–Li alloy. After immersion in 3·5 wt-% NaCl solution for 1 h (Fig. 44b ), the rare earth conversion film almost kept a uniform surface morphology and no appreciable change can be found. XPS analysis indicated that the conversion coating was composed of CeO₂, La₂O₃, MnO₂ and Mn₂O₃. Immersion tests and electrochemical measurements provided evidence that the coatings exhibit good corrosion protection on the Mg–Li alloys. ¹⁷⁰

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44.

Surface morphology of rare earth conversion film before and after immersion in 3·5 wt-% NaCl solution

Molybdate and molybdate/permanganate conversion coatings

Molybdate and molybdate/permanganate conversion coating is also a conversion coating treatment for Mg–Li alloys. ¹⁷¹ These treatments are more environmentally friendly and offer an alternative to conventional chromate conversion coatings. Figure 45 shows the surface morphology of Mg–8·5Li alloy after molybdate (Fig. 45a ) and molybdate/permanganate (Fig. 45b ) conversion treatment in different solutions. The coatings nearly covered all over the surface of the alloy, while some cracks can be seen on the coating. The molybdate/permanganate conversion coating showed compact double layer with the appearance of cracks. MoO₂, MgO, (MoO₃)_x(P₂O₅)_y and MoO₃ were the main compositions of the molybdate conversion coating. The optimum concentration of ammonium molybdate was 14 g/L. The molybdate conversion coating had better corrosion resistance. MoO₂, MgO, Mn₂O₃, MoO₃, Mn₃(PO₄)₂, MnO₂ and (MoO₃)_x(P₂O₅)_y were the main compositions of the molybdate/permanganate conversion coatings. The molybdate/permanganate conversion coating had better corrosion resistance than the molybdate conversion coating when the concentrations of permanganate and molybdate were 3·5 and 14 g/L, respectively.

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45.

Surface morphology of molybdate (a, 14 g/L) and molybdate/permanganate (b, 14 g/L molybdate+3·5 g/L permanganate) conversion coating

Stannate conversion coatings

Stannate conversion coatings have been developed to protect conventional Mg–Li alloys ^{122,172–174} since they are environmentally acceptable. Recently, stannate coating was formed on Mg–8Li alloy by the simple immersion method.

After mechanical finishing, alkaline wash, and pickling, the Mg–Li alloys samples were immersed in a stannate bath for selected periods of time. The treatment gains a 2–3 mm thick, adherent and continuous crystalline coating of MgSnO₃ on Mg–Li alloys. There was an increase in the corrosion potential of the Mg–Li alloys substrates during the formation of the stannate coatings, indicating that the coatings have a passive effect on the surface. Yang et al. reported that the coating particles were mainly composed of hemispherical particle shapes of MgSnO₃·3H₂O. ¹⁷⁵ The treatment time influences the quality of stannate conversion coatings. The stannate coatings treated for 60 min are the densest and uniform, leading to excellent corrosion resistance in their studies.

Anodic oxidation mechanism

The surface appearance of the anodised coating was as given in Fig. 46: (a) distinct rough surface with cracks and large pore; (b) uniform, cracks and less porous; (c) homogeneous, no cracks and less porous; and (d) uniform, cracks and large porous.

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46.

Surface morphology of anodic film with different deposition time: a 5 min; b 10 min; c 20 min; d 30 min

The formation of non-uniform and porous surface morphology of the film is caused by non-uniform sparking occurring on alloy surface, previous oxide on Mg–Li alloy surface melting, and gas evolution during the anodising process.

The coating layers cover the Mg–Li alloy substrate uniformly and compactly, but there were many micropores on the substrate. During the anodic oxidation process, there is gas evolution. When the maximum voltage was reached, there were sparks emerged on the surface of Mg–Li alloys, and then a large amount of heat was released. Because the heat was absorbed by the solution, the anodic oxidation film contracted. This brings about the porous structure. From a microcosmic aspect, although the thickness of the anodic oxidation film was unformed, the substrate’s chemical properties and electrochemical properties were not even. Consequently, the growth of the layer was not balanced and the distribution of the pores was also not balanced because the coatings have porous structure that could be coloured and sealed. The pore sizes are determined by the type of temperature, current density, applied voltage, electrolyte and its concentration. The pore size and density determine the quality and extent of sealing of the anodised coating. ¹⁵⁷ These coatings may be used to improve paint adhesion to the metal, as a passivation treatment or as a key for dyeing.

When the alloy is connected to a stainless steel cathode and immersed in the anodising solution, the magnesium–lithium alloys surface dissolves to a small extent in the solution, and current flows from the Mg–Li alloy to the stainless steel. The films have a thin barrier layer at the metal–coating interface followed by a layer that has a cellular structure. The electrochemistry of the coating deposition can be represented in the following equations.

Anodic reactions: Cathodic reactions:

Black anodising of a magnesium–lithium alloy

Sharma et al. ¹⁷⁶ prepared coatings in a wide range of concentrations (20–30 g/L of each constituent) and electrolyte pH (0–6·0) using the process of black anodising on Mg–Li alloy. The research results indicate that proper heat treatment could improve the microhardness of the anodised coating. The hardness of the coatings increases with increasing electrolyte concentration at first, and then decreases. The anodised coatings not only improve thermal emittance (>0·90) and solar absorptance but they can also minimise temperature gradients across the specimens. This is suitable for a particular application on the internal components of spacecraft. The pre-treatment of the black anodising of an Mg–Li alloy includes polish to obtain an even surface→flush with tap water→alkaline cleaning→flush with distilled water→acid pickling→flush with distilled water→anodising. In general, the black anodic coatings on Mg–Li alloys were obtained with the following steps:

Colouring and sealing of anodised coatings

ZSM-5 zeolite coatings onto Mg–Li alloys

Zeolites have been called excellent ‘building blocks’ for constructing hierarchical porous materials. They can be a component for functional coatings, such as low dielectric constant coatings for generating computer chips in the future. ^{164,200–202} Zeolites are a class of microporous crystalline aluminosilicates. High-silica or pure silica zeolites are especially known for their high chemical, thermal and mechanical stability. ²⁰³ ZSM-5 does not lose its crystallinity up to 1100°C. It does not react with any mineral acid except hydrofluoric acid. ²⁰⁴ ZSM-5 zeolite coatings formed by in situ crystallisation appear to offer general corrosion protection for the Mg–Li alloys. The method of hot pressing was used to assemble ZSM-5 zeolite involving the synthesis with two kinds of organic amines as templates on the surface of Mg–Li alloy. The organic amines of tetrapropylammonium bromide and tetrapropylammonium hydroxide as macromolecular templates were used to synthesise ZSM-5 zeolite. The micropores of ZSM-5 zeolite were blocked by organic templates, which enhanced the corrosion resistance of ZSM-5 coatings.

Cheng et al. ²⁰⁵ reported a novel approach to hot pressing to assemble ZSM-5 coatings on Mg–Li alloy. X-ray diffraction patterns and SEM images showed that ZSM-5 coatings are uniform and compact. ZSM-5 exhibits a pure MFI structure. The results of immersion and electrochemical tests indicated that ZSM-5 coatings improved the corrosion resistance of Mg–Li alloy because the templates block the pores of ZSM-5.

Diamond-like carbon coating on Mg–Li alloy

Diamond-like carbon (DLC) coatings on Mg–Li alloys have been investigated to improve the wear and corrosion resistance of the material. ²⁰⁶ Both deposition of a silicon (Si) interlayer and mechanical pre-treatment were applied to improve adhesion, before coating with the DLC film. ²⁰⁷

Yamauchi et al. ^208,209 studied the effectiveness of DLC coating technology in improving the wear and corrosion resistance of the Mg–14 mass-%Li alloy. The effects of two different pre-treatment methods were examined. One is coating with a Si interlayer using the ion beam sputter method and another one is peening specimens using a silicon carbide (SiC) medium before the deposition of the DLC film. The RF (Radio Frequency) capacitively coupled plasma CVD (Chemical Vapor Deposition) method was used for DLC coating in their experiments. The results indicated that the SiC peening process enhances the formation of DLC adhesive coating, leading to a relatively low wear and low friction coefficient. However, the DLC coating did not form adhesively on the substrate with a Si interlayer. Unfortunately, none of the DLC-coated specimens were able to withstand the corrosion effects of the alkaline and acidic artificial perspiration solutions.

HTMZ coating on Mg–Li alloy

A nano-sized lamellar molybdate pillared hydrotalcite/in situ created ZnO composite has been used to improve the corrosion protection for magnesium–lithium alloy. Molybdate is an excellent inhibitor of metallic alloy corrosion. ^210,211 Nano zinc oxide can polarise cathode areas by precipitating soluble salts. ^212,213 Consequently, molybdate pillared hydrotalcite/ZnO incorporated composites (HTMZ) have perfect anti-corrosion properties for Mg–Li alloy. While nano ZnO particles are an important factor controlling the corrosion inhibition by MoO₄ ²⁻ ion, HT-MoO₄ ²⁻ can release the MoO₄ ²⁻ ion and trap the corrosive Cl⁻ ion. ²¹⁴

Yu et al. ²¹⁵ has reported that nano-sized molybdate pillared hydrotalcite (HT-MoO₄ ²⁻)/ZnO incorporated composite (HTMZ) was successfully synthesised by an in situ coprecipitation method. A given amount of molybdate sodium was adjusted at pH 10·0 with NaOH and then mixed with a second solution to prepare the materials. The second solution contained Al(NO₃)₃·9H₂O and Zn(NO₃)₂·6H₂O with a metal ratio Zn²⁺/Al³⁺of 2·5. An epoxy resin coating containing HTMZ was formed on Mg–Li alloy. The research presents the comparative corrosion protection performance of the coating. Corrosion protection properties for Mg–Li alloy based on the HTMZ nanocomposite was investigated in NaCl solution by salt spray test and EIS. ²¹⁶ The results indicated that coatings containing HTMZ have better corrosion-resistant property. The mechanism of corrosion protection was attributed to the existence of ZnO, thus promoting molybdate accumulation on the Mg–Li alloy substrate.