**Influence of Solute Atoms on Deformation Behaviour of Selected Magnesium Alloys**

Zuzanka Trojanová, Peter Palček, Pavel Lukáč and Zdeněk Drozd

Additional information is available at the end of the chapter

http://dx.doi.org/10.5772/58949

### **1. Introduction**

Magnesium alloys due to their relative high specific strength and low density are used for a wide range of applications – for example in the automobile and transportation industries; they promise significant weight saving. Numerous studies have been performed in order to estimate their mechanical and physical properties. It is well known that many physical and mechanical properties of metals are influenced by alloying. The solute atoms cause an increase in the strengthening of materials. On the other hand, it is difficult to predict the effect of solute atoms on the strength and deformation behaviour of magnesium and its alloys because the experimental data concerning the critical resolved shear stress (CRSS) of single crystals of binary alloys are restricted.

In this chapter we deal with the effect of solute (foreign) atoms on the mechanical properties of magnesium. The foreign atoms are assumed to substitute for the matrix atoms and/or form precipitates. We restrict ourselves mainly to the yield stress and the true stress-true strain curves measured at a constant strain rate at room and higher temperatures. The deformation behaviour of polycrystals is influenced by the presence of crystal defects. Strength of a material is a result of strengthening mechanisms. One of the strengthening mechanisms is solution hardening defined as the increase of the initial flow stress as foreign atoms are dissolved in it [1, 2]. In this chapter original results of authors were used.

## **2. Fundamentals of plastic deformation**

Plastic deformation of magnesium polycrystals occurs by glide of dislocations and/or twin‐ ning. Glide of dislocations with the Burgers vector of 1 / 3 112 ¯0 on (0001) basal plane is the

easiest slip mechanism; often called basal slip of <*a*> type dislocation. The CRSS is required to activate the glide of <*a*> type dislocations. In order to estimate the CRSS (the stress acting on the slip plane in the slip direction) for a binary magnesium alloy (Mg-X), experimental tests on single crystals are needed. The yield strength of polycrystals is connected with the CRSS of single crystals with help of the relationship σy=M*τ*0, where M is the Taylor orientation factor. The deformation behavior of magnesium and magnesium alloy polycrystals may be influenced by twinning and the activity of non-basal slip systems. Glide of <*a*> dislocations in prismatic planes and glide of<*c*+*a*> dislocations in the second-order pyramidal slip systems should be consider. However for explanation, it is important to consider not only the solute influence but also the effect of grain size. These factors may affect not only basal slip but also prismatic slip and twinning responsible for the mechanical properties. Crystallographic textures can also change the value of the yield strength as for instance for the case of rolled sheets deformed in the rolling direction [3, 4].

It is widely accepted that the stress necessary for the dislocation motion in the slip plane can be divided into two components:

$$
\sigma\_{\rm D} = \sigma\_{\rm i} + \sigma^\*, \tag{1}
$$

where σ<sup>i</sup> is the internal (athermal) contribution to the stress, resulting from long-range internal stresses impeding the plastic flow.

$$
\sigma\_{\mathbf{i}} = a\_{\mathbf{i}} \mathbf{G} \mathbf{b} \,\rho\_{\mathbf{i}}^{1/2} \,\prime \tag{2}
$$

where G is the shear modulus, σ<sup>1</sup> is a constant describing interaction between dislocations, b is the Burgers vector of dislocations and ρ<sup>t</sup> is the total dislocation density. The effective stress σ\* acts on dislocations during their thermally activated motion when they overcome short range obstacles. The mean velocity of dislocations υ is connected with the plastic strain rate by the Orowan equation:

$$
\dot{\varepsilon} = (1/M)\,\rho\_m \, bv \tag{3}
$$

where ρ m is the mobile dislocation density. The most common equation used in describing the average dislocation velocity as a function of the effective stress is an Arrhenius type. The plastic strain rate *ε*˙ for a single thermally activated process can be expressed as:

$$\dot{\varepsilon} = \dot{\varepsilon}\_0 \exp\left[-\Delta G \left(\sigma^\*\right)/kT\right]\_{\prime} \tag{4}$$

where *ε*˙ 0 is a pre-exponential factor containing the mobile dislocation density, the average area covered by the dislocations in every activation act, the dislocation Burgers vector, the vibration frequency of the dislocation line, and the geometric factor. T is the absolute temper‐ ature and k is the Boltzmann constant. ΔG(σ\* ) is the change in the Gibbs free energy depending on the effective stress (thermal stress)σ \*=σ ap−σ<sup>i</sup> and the simple form is

$$
\Delta\!\!\Delta\!\!\!/\sigma^\*\/\!\!\/= \Delta\!\!\!\/\!G\_0\text{--}\,\text{V}\,\sigma^\* = \Delta\!\!\!\/\!G\_0\text{--}\,\text{V}\,(\sigma\cdot\sigma\_i).\tag{5}$$

Here ΔG0 is the Gibbs free energy necessary for overcoming a short range obstacle without the stress (the barrier activation energy at zero stress) and V=bdL is the activation volume where d is the obstacle wide and L is the mean length of dislocation segments between obstacles. It should be mentioned that L may depend on the stress acting on dislocation segments. In a stress relaxation (SR) test, the specimen is deformed to a certain stress σ<sup>0</sup> and then the machine is stopped and the stress is allowed to relax. The stress decreases with the time t. The specimen can be again reloaded to a higher stress (load) and the SR test may be repeated. The time derivative *σ*˙ =dσ/dt is the stress relaxation rate and σ=σ(t) is the flow stress at time t during the SR. SR tests are very often analysed under the assumption that the SR rate is proportional to the strain rate *ε*˙ and then σ(t) can be expressed according to [5] as:

$$
\Delta\sigma(\mathbf{t}) = \sigma(0) \cdot \sigma(\mathbf{t}) = a \ln(\beta \,\mathrm{t} \,\mathrm{t} \,\mathrm{t} \,\mathrm{t}) \,, \tag{6}
$$

where σ (0) ≡σ<sup>0</sup> is the stress at the beginning of the stress relaxation at time t=0, β is a constant and α=k*T*/*V*.

#### **3. Mg-Li-Al (LAxy) alloys**

easiest slip mechanism; often called basal slip of <*a*> type dislocation. The CRSS is required to activate the glide of <*a*> type dislocations. In order to estimate the CRSS (the stress acting on the slip plane in the slip direction) for a binary magnesium alloy (Mg-X), experimental tests on single crystals are needed. The yield strength of polycrystals is connected with the CRSS of single crystals with help of the relationship σy=M*τ*0, where M is the Taylor orientation factor. The deformation behavior of magnesium and magnesium alloy polycrystals may be influenced by twinning and the activity of non-basal slip systems. Glide of <*a*> dislocations in prismatic planes and glide of<*c*+*a*> dislocations in the second-order pyramidal slip systems should be consider. However for explanation, it is important to consider not only the solute influence but also the effect of grain size. These factors may affect not only basal slip but also prismatic slip and twinning responsible for the mechanical properties. Crystallographic textures can also change the value of the yield strength as for instance for the case of rolled sheets deformed in

It is widely accepted that the stress necessary for the dislocation motion in the slip plane can

s

is the internal (athermal) contribution to the stress, resulting from long-range internal

1/2

 r

where G is the shear modulus, σ<sup>1</sup> is a constant describing interaction between dislocations, b

σ\* acts on dislocations during their thermally activated motion when they overcome short range obstacles. The mean velocity of dislocations υ is connected with the plastic strain rate

(1 / ) *M bv <sup>m</sup>*

 r

where ρ m is the mobile dislocation density. The most common equation used in describing the average dislocation velocity as a function of the effective stress is an Arrhenius type. The plastic

( ) \*

 s

e

strain rate *ε*˙ for a single thermally activated process can be expressed as:

0 ee

(1)

= Gb , (2)

is the total dislocation density. The effective stress

& = (3)

= -D exp / , é ù *G kT* ê ú ë û & & (4)

D i = + \*, ss

i1 t sa

the rolling direction [3, 4].

where σ<sup>i</sup>

be divided into two components:

4 Magnesium Alloys - Properties in Solid and Liquid States

stresses impeding the plastic flow.

is the Burgers vector of dislocations and ρ<sup>t</sup>

by the Orowan equation:

#### **3.1. Microstructure of Mg-Li alloys**

Among Mg alloys, magnesium-lithium alloys, as the lightest metallic materials, are attractive for a large amount of applications. They are of great importance also for medicine purposes. Therefore, it is important to investigate mechanical properties at different temperatures and to estimate the deformation mechanisms responsible for the deformation behaviour of Mg-Li alloys at elevated temperatures. Pure Mg has hexagonal close packed structure. The density of Mg-Li alloy decreases with an increase of lithium content. The addition of Li increases ductility. The Mg–Li phase diagram shows that Li is soluble in hcp *α*-phase up to 4 wt%, while Mg alloyed with greater than 12 wt% Li has a bcc structure (β-phase) [6]. Ductility of the hcp *α*-phase are worse in comparison with the bcc alloys that are very good machinable and weldable. Disadvantages of Mg-Li alloys with bcc structure are a high chemical activity and poor corrosion resistivity. Some compromise would be an alloy with 8 wt% of Li (a mixture of phases *α*+*β*) that might exhibit both improved mechanical properties as well as a good corrosion resistance. In light micrograph in Fig. 1, light *α*-phase and darker *β*-phase may be visible. The alloys were produced by pressure infiltration under an argon pressure (up to 6 MPa) at temperatures of 615-635 K.

**Figure 1.** Light micrograph of the Mg8Li alloy.

#### **3.2. Deformation of Mg-Li alloys**

Compression tests were performed in an Instron type machine at a constant crosshead speed giving an initial strain rate of 2.4x10-4 s-1. The argon atmosphere was used as a protecting atmosphere in the furnace at elevated temperatures. The compression yield stress (CYS), σ02, was estimated as the flow stress at 0.2% offset strain. The ultimate compression strength (MCS), σmax, corresponds to the maximum value of the flow stress. Samples were deformed to a predetermined strain of 0.3, and then deformation was interrupted.

**Figure 2.** True stress-true strain curves estimated for Mg4Li alloy at various temperatures.

Influence of Solute Atoms on Deformation Behaviour of Selected Magnesium Alloys http://dx.doi.org/10.5772/58949 7

**Figure 3.** Temperature variation of the CYS and MCS estimated for Mg4Li alloy.

visible. The alloys were produced by pressure infiltration under an argon pressure (up to 6

Compression tests were performed in an Instron type machine at a constant crosshead speed giving an initial strain rate of 2.4x10-4 s-1. The argon atmosphere was used as a protecting atmosphere in the furnace at elevated temperatures. The compression yield stress (CYS), σ02, was estimated as the flow stress at 0.2% offset strain. The ultimate compression strength (MCS), σmax, corresponds to the maximum value of the flow stress. Samples were deformed to a

predetermined strain of 0.3, and then deformation was interrupted.

**Figure 2.** True stress-true strain curves estimated for Mg4Li alloy at various temperatures.

MPa) at temperatures of 615-635 K.

6 Magnesium Alloys - Properties in Solid and Liquid States

**Figure 1.** Light micrograph of the Mg8Li alloy.

**3.2. Deformation of Mg-Li alloys**

**Figure 4.** True stress-true strain curves estimated for Mg8Li alloy at various temperatures.

Figure 2 shows the true stress-true strain curves estimated at various temperatures. A significant hardening is obvious especially for lower temperatures. Figure 3 shows the temperature variation of the CYS, σ02, as well as the MCS (σmax) of Mg4Li alloy. It is apparent from Fig. 3 that the temperature variation of CYS exhibits a local maximum at a temperature of 70 °C. The MCS of Mg-4Li alloy decreases rapidly with increasing temperature. The differences between MCS and CYS exhibit a rapid decrease with increasing temperature. The true stress-true strain curves of Mg-8Li alloy measured in compression at various temperatures are shown in Fig. 4. A rapid decrease of the flow stresses at temperatures higher than 50 °C is obvious in Fig. 5. The temperature variation of the CYS exhibits a local maximum at 50 °C. While the stress-strain curves estimated for Mg4Li exhibit a significant strain hardening, the curves observed for Mg12Li are very flat (Fig. 6); the difference between the CYS and MCS is relatively low and decreases with increasing temperature (Fig. 7).

The Mg-4Li alloy has *hcp* structure. The addition of Li to Mg increases the critical resolved shear stress (CRSS) for basal slip; the solid solution hardening is observed [7]. The CRSS of Mg4Li is almost independent on the temperature above room temperature while the CRSS for non-basal slip decreases with temperature below 200 °C as shown in [8]. The addition of Li causes a decrease of both *a* and *c* lattice parameters in Mg-Li solid solution [9].

**Figure 5.** Temperature variation of the CYS and MCS estimated for Mg8Li alloy.

**Figure 6.** True stress-true strain curves of Mg12Li alloy at various temperatures.

It is well known that the dominant slip system in Mg and *hcp* Mg alloys at room temperature is the basal one. To fulfil von Mises criterion, a non-basal slip system should be active. The activity of non-basal slip systems plays an important role in dynamic recovery (softening). The pyramidal slip systems can be considered as non-basal slip systems. During deformation of magnesium alloy polycrystals, the motion not only <**a>** (basal) dislocations but also **<***c***+***a***>** (pyramidal) dislocations is assumed. Screw components of the <*c*+*a*> dislocations can move to the parallel slip planes by double cross slip and then annihilate, which causes a decrease in the strain hardening rate; softening is observed. With the addition of Li the c/a ratio decreases, which may result in a higher activity of non-basal slip [10]. Agnew et al. [11] revealed that the <*c*+*a*> dislocations in the pyramidal planes improve ductility of MgLi alloys. Pawelek et al. [12] studying acoustic emission from deformed Mg-Li alloys estimated a high level of acoustic emission in Mg4Li alloy as a result of non-basal slip in the prismatic and pyramidal slip systems.

**Figure 7.** Temperature variation of the CYS and MCS estimated for Mg12Li alloy.

#### **3.3. Microstructure of Mg-xLi-yAl alloys**

While the stress-strain curves estimated for Mg4Li exhibit a significant strain hardening, the curves observed for Mg12Li are very flat (Fig. 6); the difference between the CYS and MCS is

The Mg-4Li alloy has *hcp* structure. The addition of Li to Mg increases the critical resolved shear stress (CRSS) for basal slip; the solid solution hardening is observed [7]. The CRSS of Mg4Li is almost independent on the temperature above room temperature while the CRSS for non-basal slip decreases with temperature below 200 °C as shown in [8]. The addition of Li

causes a decrease of both *a* and *c* lattice parameters in Mg-Li solid solution [9].

relatively low and decreases with increasing temperature (Fig. 7).

8 Magnesium Alloys - Properties in Solid and Liquid States

**Figure 5.** Temperature variation of the CYS and MCS estimated for Mg8Li alloy.

**Figure 6.** True stress-true strain curves of Mg12Li alloy at various temperatures.

To improve the mechanical properties, alloying with other elements can be used. The addition of Al atoms to Mg4Li causes the formation of precipitates. A combination of two different hardening mechanisms should be considered: solid solution hardening and precipitation hardening. The microstructure of as-cast Mg8LixAl alloys under consideration consists essentially of *hcp* magnesium-based *α*-phase and *bcc* lithium-based one (*β*-phase) and AlLi precipitates as well. Accordingly, interpenetrated (α+β) matrix structure with dominating αphase is characteristic for LA83 and LA85 (see Fig. 8) alloys. Light micrograph of the Mg12Li3Al (LA123) alloy shows that the alloy does not contain only one phase (Fig. 9). Both phases (*α* and *β*) are present contrary to the alloy without Al because aluminium stabilises the *hcpα*-phase.

**Figure 8.** Light micrograph of the Mg8Li5Al alloy.

**Figure 9.** Light micrograph of the Mg12Li3Al.

**Figure 10.** Temperature variation of the CYS and MCS estimated for Mg4Li3Al alloy.

#### **3.4. Deformation of Mg-xLi-yAl alloys**

**Figure 8.** Light micrograph of the Mg8Li5Al alloy.

10 Magnesium Alloys - Properties in Solid and Liquid States

**Figure 9.** Light micrograph of the Mg12Li3Al.

**Figure 10.** Temperature variation of the CYS and MCS estimated for Mg4Li3Al alloy.

The temperature dependence of the CYS and MCS estimated for LA43 (Mg-4Li-3Al) and LA45 (Mg-4Li-5Al) are introduced in Figs. 10 and 11. High differences between CYS and MCS, especially observed at lower temperatures, and moderate decrease of the characteristic stresses with temperature indicate significant hardening during plastic deformation. The presence of AlLi precipitates influences not only the yield stress but also the storage of dislocations during plastic deformation. In the LA43 (Mg4Li3Al) and LA45 (Mg4Li5Al) alloys, there are obstacles of non-dislocation types such as precipitates and the dislocation obstacles (forest dislocations). The observed high difference between the CYS and MCS indicates a significant hardening during the deformation process. A part of the moving dislocations stored at the obstacles contributes to hardening. On the other hand, processes such as cross slip and climb of dislocations contribute to softening– the difference between CYS and MCS decreases with increasing deformation temperature. The dislocation microstructure can change. For simplic‐ ity, the total dislocation density is considered as the characteristic parameter of the evolution of microstructure during deformation. According to the model of Lukáč and Balík [13], we take into account storage of dislocations at both impenetrable obstacles and forest dislocations, and annihilation of dislocations due to both cross slip and climb.

**Figure 11.** Temperature variation of the CYS and MCS estimated for Mg4Li5Al alloy.

The true stress-true strain curves of LA85 alloy estimated at various temperatures are shown in Fig. 12. The strain hardening is observed at temperatures up to 100 °C; the stress-strain curves obtained at higher temperatures are flat. The temperature variations of the CYS and MCS for LA83 and LA85 alloys are shown in Figs. 13 and 14. The strengthening effect of Al atoms in the case of LA85 alloy has been found to be really high even at 100 °C. A moderate decrease of both characteristic stresses with temperature classes this alloy as a material for industrial applications.

**Figure 12.** True stress-true strain curves of Mg8Li5Al alloy at various temperatures.

**Figure 13.** Temperature variation of the CYS and MCS estimated for Mg8Li3Al alloy.

**Figure 14.** Temperature variation of the CYS and MCS for Mg8Li5Al alloy.

Influence of Solute Atoms on Deformation Behaviour of Selected Magnesium Alloys http://dx.doi.org/10.5772/58949 13

**Figure 15.** Temperature variation of the CYS and MCS estimated for Mg12Li3Al.

**Figure 12.** True stress-true strain curves of Mg8Li5Al alloy at various temperatures.

12 Magnesium Alloys - Properties in Solid and Liquid States

**Figure 13.** Temperature variation of the CYS and MCS estimated for Mg8Li3Al alloy.

**Figure 14.** Temperature variation of the CYS and MCS for Mg8Li5Al alloy.

**Figure 16.** Temperature variation of the CYS and MCS estimated for Mg12Li5Al.

The temperature variations of both CYS and MCS are shown for LA123 and LA125 alloys in Figs. 15 and 16, respectively. The CYS is significantly higher for the alloy containing 5% of Al, while MCS exhibits higher values for the LA123 alloy. Thermal stability is higher in the case of LA125 alloy. The influence of Al on the mechanical characteristics of Mg-Li alloys at two temperatures is summarised in Table 1.


**Table 1.** Influence of Al addition on CYS and MCS of Mg4Li, Mg8Li and Mg12Li alloys.

To conclude it is possible to say that the best mechanical properties of the superlight MgLi alloys were found for Mg8Li alloy with 5 wt% of Al. Combining of the stronger *α* phase with a more ductile *β* phase results in a material with the high specific strength.

#### **4. Mg-Al-Sr (AJxy) alloys**

#### **4.1. Microstructure of the Mg-Al-Sr alloys**

Special industrial applications require improvement of the high temperature properties. For these elevated temperature applications, alloys containing rare earth elements have been developed. New Mg-Al-Sr alloys are being developed with the aim to find cast alloys with good creep resistance and good strength and replace expensive rare earth alloying elements with some cheaper one. Pekguleryuz [14] reported that Mg-Al-Sr alloys show different microstructures based on the Sr/Al ratio. For Sr/Al ratio below about 0.3, only Al4Sr interme‐ tallic phase is present as the second phase in the structure. When the Sr/Al ratio is higher, a second intermetallic phase, a new, ternary Mg-Al-Sr compound, is observed. When the Sr/Al ratio is very low, there is insufficient amount of Sr to bind all Al atoms and the excess Al would form the Mg17Al12 phase. Figure 17 shows light micrograph of the squeeze cast AJ51 alloy. The primary Mg grains are surrounded by the interconnected network of the grain boundary phase. This phase is formed during solidification process and it has lamellar type morphology. The *γ* phase (Mg17Al12) is accompanied with Al4Sr phase. Single items are better visible in Figs. 18 and 19, where SEM of AJ91 alloy is presented. Eutectic *γ* phase forms either continuous shapes or discontinuous precipitates. Light needles visible in both pictures are the Al4Sr precipitates. Line scan of alloying elements present in particles shown in Fig. 19 is given in Fig. 20. Line scan analysis was performed along the light arrow showed in Fig. 19. Small content of Sr is present in the *δ* solid solution; higher in the position of a needle shaped particle.

**Figure 17.** Light micrograph of the AJ51 alloy.

#### **4.2. Deformation of the AJ alloys**

To conclude it is possible to say that the best mechanical properties of the superlight MgLi alloys were found for Mg8Li alloy with 5 wt% of Al. Combining of the stronger *α* phase with

Special industrial applications require improvement of the high temperature properties. For these elevated temperature applications, alloys containing rare earth elements have been developed. New Mg-Al-Sr alloys are being developed with the aim to find cast alloys with good creep resistance and good strength and replace expensive rare earth alloying elements with some cheaper one. Pekguleryuz [14] reported that Mg-Al-Sr alloys show different microstructures based on the Sr/Al ratio. For Sr/Al ratio below about 0.3, only Al4Sr interme‐ tallic phase is present as the second phase in the structure. When the Sr/Al ratio is higher, a second intermetallic phase, a new, ternary Mg-Al-Sr compound, is observed. When the Sr/Al ratio is very low, there is insufficient amount of Sr to bind all Al atoms and the excess Al would form the Mg17Al12 phase. Figure 17 shows light micrograph of the squeeze cast AJ51 alloy. The primary Mg grains are surrounded by the interconnected network of the grain boundary phase. This phase is formed during solidification process and it has lamellar type morphology. The *γ* phase (Mg17Al12) is accompanied with Al4Sr phase. Single items are better visible in Figs. 18 and 19, where SEM of AJ91 alloy is presented. Eutectic *γ* phase forms either continuous shapes or discontinuous precipitates. Light needles visible in both pictures are the Al4Sr precipitates. Line scan of alloying elements present in particles shown in Fig. 19 is given in Fig. 20. Line scan analysis was performed along the light arrow showed in Fig. 19. Small content of Sr is present in the *δ* solid solution; higher in the position of a needle shaped particle.

a more ductile *β* phase results in a material with the high specific strength.

**4. Mg-Al-Sr (AJxy) alloys**

**Figure 17.** Light micrograph of the AJ51 alloy.

**4.1. Microstructure of the Mg-Al-Sr alloys**

14 Magnesium Alloys - Properties in Solid and Liquid States

The true stress-true strain curves of AJ51 alloy deformed in compression at various tempera‐ tures are shown in Fig. 21. Samples were deformed either to failure or the tests performed at temperatures higher than 100 °C were interrupted at a predetermined strain. Significant hardening at temperatures up to 150 °C is obvious. The curves obtained at temperatures higher are more or less flat – hardening and softening processes are in equilibrium.

**Figure 18.** Scanning electron micrograph of as cast AJ91 alloy.

**Figure 19.** Scanning electron micrograph of as cast AJ91 alloy (different place).

**Figure 20.** Line scan of alloying elements.

Figure 22 shows the temperature variation of the CYS as well as the MCS. Similar character‐ istics estimated in tension are shown in Fig. 23. While the values of the yield stress are practically the same for tension and compression, the values of the maximum stress are higher in compression tests. It is a consequence of higher ductility of the alloy in compression and significant hardening during plastic deformation at lower temperatures. The yield stress decreases with increasing temperature for samples deformed in compression. A small local maximum in the temperature dependence of the yield strength observed in the vicinity of 50 °C was observed in several Mg alloys and composites [15]. It is likely caused by a dynamic age hardening.

**Figure 21.** True stress-true strain curves estimated in compression for AJ51 alloy at various temperatures.

The true stress-true strain curves estimated for AJ91 alloy in compression are presented in Fig. 24. The values of the CYS and TYS (Figs. 25 and 26) are higher than the yield stress of the mostly used cast alloy AZ91. The temperature variations of MCS and MTS are shown in Figs. 25 and 26. Ductility of AJ91 alloy deformed in tension is low, at ambient temperature only several percent. Thermal stability is in comparison with AZ91 alloy also better; the CYS as well as TYS do not decrease with increasing test temperature up to 200 °C below 100 MPa. It is done by the thermal stability of Al4Sr precipitates. Solid solution hardening plays in this case only minor role.

**Figure 22.** Temperature variation of CYS and MCS for AJ51 alloy.

**Figure 20.** Line scan of alloying elements.

16 Magnesium Alloys - Properties in Solid and Liquid States

hardening.

Figure 22 shows the temperature variation of the CYS as well as the MCS. Similar character‐ istics estimated in tension are shown in Fig. 23. While the values of the yield stress are practically the same for tension and compression, the values of the maximum stress are higher in compression tests. It is a consequence of higher ductility of the alloy in compression and significant hardening during plastic deformation at lower temperatures. The yield stress decreases with increasing temperature for samples deformed in compression. A small local maximum in the temperature dependence of the yield strength observed in the vicinity of 50 °C was observed in several Mg alloys and composites [15]. It is likely caused by a dynamic age

**Figure 21.** True stress-true strain curves estimated in compression for AJ51 alloy at various temperatures.

The true stress-true strain curves estimated for AJ91 alloy in compression are presented in Fig. 24. The values of the CYS and TYS (Figs. 25 and 26) are higher than the yield stress of the mostly used cast alloy AZ91. The temperature variations of MCS and MTS are shown in Figs. 25 and

**Figure 23.** Temperature variation of TYS and MTS for AJ51 alloy.

#### **4.3. Stress components in the AJ alloys**

Stress necessary for dislocation motion is possible, according to eq. (1), divided into two components. The components of the applied stress (σ<sup>i</sup> , σ\* ) may be estimated using Li's method [16]. The SR curves are fitted to the power law function in the form:

$$
\sigma - \sigma\_i = \left[ a \left( m - 1 \right) \right]^{\frac{1}{1 - m}} \left( t + t\_0 \right)^{\frac{1}{1 - m}} \tag{7}
$$

where a, t0 and m are fitting parameters. The influence of solute atoms on both stress compo‐ nents of AJ51 and AJ91 is obvious from Figs 27-29; it was estimate at three temperatures. At room temperature, an increase in the concentration of solute atoms influences only the internal stress. The observed increase of the internal stress is due to higher density of impenetrable precipitates which are obstacles for the dislocation motion. The effective stress σ\* is practical‐ ly the same for both alloys. From Fig. 27 it is obvious that the internal stress in both alloys is extremely high; it represents more than 90% of the applied stress. The difference between the internal stress of AJ51 and AJ91 alloys at 100 °C (Fig. 28) is relatively high and it indicates the reinforcing effect of the increased concentration of solute atoms. The effective stress is still for both alloys practically the same but the values of σ\* for both alloys are higher in comparison with the values obtained at room temperature. The internal stress estimated for AJ91 alloy at 200 °C (Fig. 29) is for strains up to approximately ε~ 0.08 higher as the effective stress. While the internal stress for AJ91 alloy decreases with strain, the effective stress continuously increases for both alloys. The decrease of the internal stress with strain estimated for both alloys is a consequence of the intensive activity of softening processes. This mechanism will be dis‐ cussed in details later.

**Figure 24.** True stress-true strain curves of AJ91alloy deformed in compression at various temperatures.

**Figure 25.** Temperature variation of CYS and MCS for AJ91 alloy.

Influence of Solute Atoms on Deformation Behaviour of Selected Magnesium Alloys http://dx.doi.org/10.5772/58949 19

**Figure 26.** Temperature variation of TYS and MTS for AJ91 alloy.

where a, t0 and m are fitting parameters. The influence of solute atoms on both stress compo‐ nents of AJ51 and AJ91 is obvious from Figs 27-29; it was estimate at three temperatures. At room temperature, an increase in the concentration of solute atoms influences only the internal stress. The observed increase of the internal stress is due to higher density of impenetrable precipitates which are obstacles for the dislocation motion. The effective stress σ\* is practical‐ ly the same for both alloys. From Fig. 27 it is obvious that the internal stress in both alloys is extremely high; it represents more than 90% of the applied stress. The difference between the internal stress of AJ51 and AJ91 alloys at 100 °C (Fig. 28) is relatively high and it indicates the reinforcing effect of the increased concentration of solute atoms. The effective stress is still for

with the values obtained at room temperature. The internal stress estimated for AJ91 alloy at 200 °C (Fig. 29) is for strains up to approximately ε~ 0.08 higher as the effective stress. While the internal stress for AJ91 alloy decreases with strain, the effective stress continuously increases for both alloys. The decrease of the internal stress with strain estimated for both alloys is a consequence of the intensive activity of softening processes. This mechanism will be dis‐

**Figure 24.** True stress-true strain curves of AJ91alloy deformed in compression at various temperatures.

for both alloys are higher in comparison

both alloys practically the same but the values of σ\*

18 Magnesium Alloys - Properties in Solid and Liquid States

**Figure 25.** Temperature variation of CYS and MCS for AJ91 alloy.

cussed in details later.

**Figure 27.** Strain dependence of the internal stress, σ<sup>i</sup> , and effective stress, σ\*, for AJ51 and AJ91 alloys deformed at 25 °C

**Figure 28.** Strain dependence of the internal stress, σ<sup>i</sup> , and effective stress, σ\*, for AJ51 and AJ91 alloys at 100 °C.

**Figure 29.** Strain dependence of the internal stress, σ<sup>i</sup> , and effective stress, σ\*, for AJ51 and AJ91 alloys at 200 °C.

#### **5. Mg-Al-Ca (AXxy) alloys**

#### **5.1. Microstructure of the AX alloys**

When Ca is added to Mg-Al binary alloys, the type of precipitating compound depends on the Ca/Al mass ratio. When this ratio is higher than 0.8 the presence of both Mg2Ca and Al2Ca were detected. They may contribute to a considerable increase in hardness and the yield stress [17]. For the ratio below 0.8 only Al2Ca Laves phase (C15-cubic) was observed to have been formed. Both types of precipitates were observed to form along the grain boundaries [18]. Gjestland et al. [19] showed that the creep resistance of AX alloy at 150 °C is similar to magnesium alloys containing rare earths with the added benefit of good corrosion resistance. Terada et al. [20] studied the creep mechanisms in the Mg-5Al-1.7Ca alloy. They found a change of deformation mechanism at the vicinity of 150 °C. Microstructure and mechanical properties of Mg-Al based alloy with Ca addition (AX series) were investigated by Du et al. [21]. They estimated that the small amount of Ca increased the thermal stability of Mg17Al12 intermetallic phase, so that the creep resistance at elevated temperatures was improved.

Microstructure of the squeeze cast AX41 alloy used in this study exhibits typical dendrit‐ ic structure with *γ* Mg grains decorated by particles. SEM showed the details of the particles structure (Fig. 30); Mg17Al12 intermetallic phase surrounded with smaller particles of Al2Ca. Microstructure of the squeeze cast AX91 alloy is displayed in Fig. 31. Two types of particles were identified, which is better evident in the back scattered electrons image (see Fig. 32). Dark particles in Fig. 32 are eutectics Mg17Al12; lighter skeleton-like particles containing Ca atoms.

**Figure 30.** Scanning electron micrograph of as cast AX41 alloy showing the main phases present in the alloy.

**Figure 29.** Strain dependence of the internal stress, σ<sup>i</sup>

20 Magnesium Alloys - Properties in Solid and Liquid States

**5. Mg-Al-Ca (AXxy) alloys**

atoms.

**5.1. Microstructure of the AX alloys**

, and effective stress, σ\*, for AJ51 and AJ91 alloys at 200 °C.

When Ca is added to Mg-Al binary alloys, the type of precipitating compound depends on the Ca/Al mass ratio. When this ratio is higher than 0.8 the presence of both Mg2Ca and Al2Ca were detected. They may contribute to a considerable increase in hardness and the yield stress [17]. For the ratio below 0.8 only Al2Ca Laves phase (C15-cubic) was observed to have been formed. Both types of precipitates were observed to form along the grain boundaries [18]. Gjestland et al. [19] showed that the creep resistance of AX alloy at 150 °C is similar to magnesium alloys containing rare earths with the added benefit of good corrosion resistance. Terada et al. [20] studied the creep mechanisms in the Mg-5Al-1.7Ca alloy. They found a change of deformation mechanism at the vicinity of 150 °C. Microstructure and mechanical properties of Mg-Al based alloy with Ca addition (AX series) were investigated by Du et al. [21]. They estimated that the small amount of Ca increased the thermal stability of Mg17Al12 intermetallic

Microstructure of the squeeze cast AX41 alloy used in this study exhibits typical dendrit‐ ic structure with *γ* Mg grains decorated by particles. SEM showed the details of the particles structure (Fig. 30); Mg17Al12 intermetallic phase surrounded with smaller particles of Al2Ca. Microstructure of the squeeze cast AX91 alloy is displayed in Fig. 31. Two types of particles were identified, which is better evident in the back scattered electrons image (see Fig. 32). Dark particles in Fig. 32 are eutectics Mg17Al12; lighter skeleton-like particles containing Ca

phase, so that the creep resistance at elevated temperatures was improved.

**Figure 31.** Scanning electron micrograph showing two main phases present in the AX91 alloy. The skeleton-like phase contains Ca.

**Figure 32.** Back scattered electrons image of the same place as showed in Fig. 31.

**Figure 33.** True stress-true strain curves of AX41 alloy deformed in tension at various temperatures.

Influence of Solute Atoms on Deformation Behaviour of Selected Magnesium Alloys http://dx.doi.org/10.5772/58949 23

**Figure 34.** Temperature variation of TYS and MTS for AX41 alloy.

**Figure 35.** Temperature variation of CYS and MCS for AX41 alloy.

#### **5.2. Deformation of the AX alloys**

**Figure 32.** Back scattered electrons image of the same place as showed in Fig. 31.

22 Magnesium Alloys - Properties in Solid and Liquid States

**Figure 33.** True stress-true strain curves of AX41 alloy deformed in tension at various temperatures.

The true stress-true strain curves of AX41 alloy deformed in tension are presented in Fig. 33. Low ductility (about 5%) of the alloy was observed at lower temperatures below 100 °C; at higher temperatures ductility increases up to 27%. The temperature variations of the TYS and MTS are presented in Fig. 34.

**Figure 36.** Temperature dependence of TYS and MTS for AX61 alloy.

**Figure 37.** True stress-true strain curves of AX91 alloy deformed in tension at various temperatures.

**Figure 38.** Temperature variation of TYS and MTS for AX91 alloy.

Analogously the values of the CYS and MCS at different temperatures are shown in Fig. 35. It can be seen that the TYS decreases with increasing temperature monotonously whereas the temperature variation of the CYS exhibits a local maximum at about 50 °C. The temperature variations of the TYS and MTS for squeeze cast AX61 alloy are presented in Fig. 36. The values of the TYS are relatively high and decrease only slightly with increasing temperature exhibiting at 200 °C still a value about 100 MPa. The true stress-true strain curves of AX91 alloy measured in tension at various temperatures are shown in Fig. 37. It is obvious from Fig. 37 that ductility of the alloy at temperatures up to 100 °C is limited; it increases up to 25 % at a temperature of 300 °C. The TYS at room temperature was estimated to be approximately equal to that measured for AX61 alloy. The observed rapid decrease of TYS of AX91 with temperature indicates a lower thermal stability of this alloy (Fig. 38). It is a different situation in comparison to the AJ91 alloy, for which the strength was found to be superior.

**Figure 39.** Strain dependence of the internal (σ<sup>i</sup> ) and effective (σ\*) stresses at 25 °C.

**Figure 36.** Temperature dependence of TYS and MTS for AX61 alloy.

24 Magnesium Alloys - Properties in Solid and Liquid States

**Figure 38.** Temperature variation of TYS and MTS for AX91 alloy.

**Figure 37.** True stress-true strain curves of AX91 alloy deformed in tension at various temperatures.

**Figure 40.** Strain dependence of the internal (σ<sup>i</sup> ) and effective (σ\*) stresses at 100 °C.

**Figure 41.** Strain dependence of the internal (σ<sup>i</sup> ) and effective (σ\*) stresses at 200 °C.

#### **5.3. Stress components in the AX alloys**

The applied stress components, σ<sup>i</sup> , and, σ\*, for AX41 and AX91 alloys were estimated at three temperatures in compression (see Figs. 39-41). At lower temperatures, 25 and 100 °C, the solute atom concentration influences only the internal stress. The effective stress is for both alloys the same (at 25°C) or it is a bit higher (at 100 °C) for AJ91 alloy. In both cases the effective stress slightly increases with increasing strain. A different behaviour was found at 200 °C. While the internal stress estimated for AX41 alloy increases with increasing strain up to 12%, the internal stress in AX91 alloy slightly increases with strain up to 8% and then decreases. On the other hand, the effective stress increases in the whole strain range.

According to eq. (2) the internal stress, σ<sup>i</sup> is proportional to ρ1/2 where ρ is the density of dislocations. The internal stress, σ<sup>i</sup> , generally, reflects the resistance of a metallic material against plastic deformation. Considering a constant microstructure, the deformation (flow) stresses are done by the evolution of the dislocation density with strain and temperature. The observed increase of the internal stress for all alloys AJ and AX series indicates an increase in the dislocation density. The moving dislocations can be stored at both non-dislocation and dislocation type obstacles. Non-dislocation obstacles may be grain boundaries, non-coherent precipitates and/or twins; the dislocation type obstacles are formed by reactions between dislocations. As mentioned in paragraph 2, in *hcp* magnesium alloys dislocations with the Burgers vector **<***a***>**=1 / 3 112 ¯0 may glide in basal, prismatic, and first−order pyramidal planes. Dislocations with the Burgers vector <*c+a*>=1 / 3 112 ¯3 can glide in the first− and second−order pyramidal planes. The basal <*a*> dislocations may react with the pyramidal <c+*a*>dislocations according to:

$$
\frac{1}{3}\langle 2\overline{1}\overline{1}3 \rangle + \frac{1}{3}\langle \overline{2}110 \rangle \to \{0001\}.\tag{8}
$$

The resulting sessile**<***c***>** dislocations with the Burgers vector parallel to the c axis are not able to glide in the basal plane; therefore such dislocations are obstacles for moving dislocations.

Twins are another type of obstacles. Twinning plays an important role in plastic deformation of *hcp* magnesium alloys. Our experiments were performed in compression. It was shown [22-24] that {101 ¯0} compressive twinning, {101 ¯1}-{101 ¯2} double twining and {101 ¯3}-{101 ¯2} double twinning could also accommodate compressive strains along the c-axis at lower temperatures. The nearly constant level or slightly decreasing tendency of the internal stress estimated at a temperature of 200 °C indicates a decrease in the dislocation density as a consequence of recovery process/-es. With rising temperature, the intensity of dynamic recovery is increasing – this can be related to dislocation climb and also to the activity of additional non-basal slip systems. It should be mentioned that <*c*> dislocations are not able to glide (conservative movement) in the basal plane, however they may climb at elevated temperatures and release the primary glide. Another dislocation reaction may yield a sessile <*c***+***a***>** dislocation:

**Figure 41.** Strain dependence of the internal (σ<sup>i</sup>

26 Magnesium Alloys - Properties in Solid and Liquid States

The applied stress components, σ<sup>i</sup>

**5.3. Stress components in the AX alloys**

According to eq. (2) the internal stress, σ<sup>i</sup>

Dislocations with the Burgers vector <*c+a*>=1 / 3 112

1 1

dislocations. The internal stress, σ<sup>i</sup>

Burgers vector **<***a***>**=1 / 3 112

according to:

hand, the effective stress increases in the whole strain range.

) and effective (σ\*) stresses at 200 °C.

temperatures in compression (see Figs. 39-41). At lower temperatures, 25 and 100 °C, the solute atom concentration influences only the internal stress. The effective stress is for both alloys the same (at 25°C) or it is a bit higher (at 100 °C) for AJ91 alloy. In both cases the effective stress slightly increases with increasing strain. A different behaviour was found at 200 °C. While the internal stress estimated for AX41 alloy increases with increasing strain up to 12%, the internal stress in AX91 alloy slightly increases with strain up to 8% and then decreases. On the other

against plastic deformation. Considering a constant microstructure, the deformation (flow) stresses are done by the evolution of the dislocation density with strain and temperature. The observed increase of the internal stress for all alloys AJ and AX series indicates an increase in the dislocation density. The moving dislocations can be stored at both non-dislocation and dislocation type obstacles. Non-dislocation obstacles may be grain boundaries, non-coherent precipitates and/or twins; the dislocation type obstacles are formed by reactions between dislocations. As mentioned in paragraph 2, in *hcp* magnesium alloys dislocations with the

pyramidal planes. The basal <*a*> dislocations may react with the pyramidal <c+*a*>dislocations

, and, σ\*, for AX41 and AX91 alloys were estimated at three

is proportional to ρ1/2 where ρ is the density of

¯3 can glide in the first− and second−order

, generally, reflects the resistance of a metallic material

¯0 may glide in basal, prismatic, and first−order pyramidal planes.

3 3 2113 2110 0001 . + ® (8)

$$
\frac{1}{3}\{2\overline{1}\overline{1}3\} + \frac{1}{3}\{\overline{1}2\overline{1}0\} \to \frac{1}{3}\{11\overline{2}3\}.\tag{9}
$$

A combination of two glissile<*c***+***a***>** dislocations gives rise to a sessile dislocation of <*a*> type that lays along the intersection of the second order pyramidal planes according to the following reaction:

$$
\frac{1}{3}\left\langle 2\overline{1}\overline{1}3 \right\rangle + \frac{1}{3}\left\langle \overline{1}2\overline{1}\overline{3} \right\rangle \to \frac{1}{3}\left\langle 11\overline{2}0 \right\rangle. \tag{10}
$$

It can be seen that different dislocation reactions may produce both sessile and glissile dislocations. Production of sessile dislocations increases the density of the forest dislocations that are obstacles for moving dislocations. Therefore, an increase in the flow stress with straining (i.e. hardening) follows, which is observed in the experiment.

Dislocations may be stored in front of impenetrable, thermally stable, precipitates and therefore dislocation pile-ups can be formed. These pile-ups are very effective stress concen‐ trators. Local stress produced by the dislocation pile-ups may support cross slip of screw dislocations and so contribute to softening of the alloy. A higher density of precipitates in AJ91 and AX91 alloys compared with AJ51 and AJ41 alloys is the main reason for higher values of the CYS/TYS observed at lower temperatures. On the other hand, the precipitates (significant stress concentrators) may make easier climb of dislocations at elevated temperatures. Higher mobility of dislocations in prismatic and pyramidal slip planes at elevated temperatures increases the probability of dislocation reactions between <*c***+***a***>** dislocations leading to dynamic recovery in a *hcp* structure.

#### **6. Influence of solute atoms on activation volume**

While the internal stress is strongly influenced with the content of solute atoms, the effective stress was– up to 100 °C–not affected by solute atoms (in the case of AJxy and AXxy alloys). The observed increase of the effective stress at higher temperatures is not surprising; the thermally activated process at higher temperatures is complex. Solute atoms (or their small clusters) are considered as typical local obstacles for moving dislocations. In high-temperature regime, diffusion-controlled glide should be taken into account.

**Figure 42.** The plot of the activation volume in b3 against the effective stress σ\* for AJ51 alloy deformed in tension and compression at various temperatures.

**Figure 43.** The plot of the activation volume in b3 against the effective stress σ\* for AJ91 alloy deformed in tension and compression at various temperatures.

**6. Influence of solute atoms on activation volume**

28 Magnesium Alloys - Properties in Solid and Liquid States

regime, diffusion-controlled glide should be taken into account.

V/b3

V/b3

compression at various temperatures.

compression at various temperatures.

While the internal stress is strongly influenced with the content of solute atoms, the effective stress was– up to 100 °C–not affected by solute atoms (in the case of AJxy and AXxy alloys). The observed increase of the effective stress at higher temperatures is not surprising; the thermally activated process at higher temperatures is complex. Solute atoms (or their small clusters) are considered as typical local obstacles for moving dislocations. In high-temperature

AJ51

T 25°C T 100°C T 150°C T 200°C C 25°C C 100°C C 200°C

C 25°C C 50°C C 100°C C 150°C C 200°C C 250°C T 200°C T 250°C

s\* (MPa) 0 20 40 60 80 100 120

**Figure 42.** The plot of the activation volume in b3 against the effective stress σ\* for AJ51 alloy deformed in tension and

AJ91

s\* (MPa) 0 20 40 60 80 100 120 140

**Figure 43.** The plot of the activation volume in b3 against the effective stress σ\* for AJ91 alloy deformed in tension and

**Figure 44.** The plot of the activation volume in b3 against the effective stress σ\* for AJ51 and AJ91 alloys in tension and compression at all temperatures.

The values of the activation volume, V, were estimated in the stress relaxation experiments using equation (6). As usual, the values of the activation volume divided by b3 for samples of AJ51 and AJ91 alloys deformed in tension as well as compression are plotted against the effective stress σ\* for all testing temperatures in Figs 42-43 for tension (empty characters) and compression (full characters) deformation. The same analysis was performed for AX41 and AX91 alloys as it can be seen in Figs. 45-47. Plotting values of V against the effective stress for both AJ alloys into one diagram (Fig. 44) and those for both AX alloys in Fig. 47 shows that the activation volumes decrease with the effective stress and all the values lie on one line –"master curve". Kocks et al. [25] suggested an empirical equation between Gibbs free energy ΔG and the effective stress σ\* in the following form:

$$
\Delta G = \Delta G\_0 \left[ 1 - \left( \frac{\sigma^\*}{\sigma\_0^\*} \right)^p \right]^q \,\tag{11}
$$

where Δ G0 and *σ*<sup>0</sup> <sup>∗</sup> are Gibbs energy and the effective stress at 0 K. For the effective stress it follows:

$$
\sigma^\* = \sigma\_0^\* \left[ 1 - \left( \frac{kT}{\Delta G\_0} \ln \frac{\dot{\varepsilon}\_0}{\dot{\varepsilon}} \right)^{1/q} \right]^{1/p} \tag{12}
$$

**Figure 45.** The plot of the activation volume in b3 against the effective stress σ\* estimated for AX41 alloy at various deformation temperatures in tension and compression.

**Figure 46.** The plot of the activation volume in b3 against the effective stress σ\* estimated for AX91 alloy at various deformation temperatures in tension and compression.

3

**Figure 47.** The plot of the activation volume in b3 against the effective stress σ\* estimated for AX41 and AX91 alloys at all temperatures in tension and compression.

where p and q are phenomenological parameters reflecting the shape of a resistance obstacle profile. The possible ranges of values p and q are limited by the conditions 0 < p≤ 1 and 1 ≤ q≤ 2. Ono [26], suggested that Equation (12) with p=1/2, q=3/2 describes a barrier shape profile that fits many predicted barrier shapes. Thermodynamics generally defines the activation volume as

$$V = -\left(\frac{d\Delta G}{d\sigma^\*}\right)\_T.\tag{13}$$

Equation (13) can be rewritten as

AX41

V/b3

V/b3

V/b3

0

100

200

300

**Figure 45.** The plot of the activation volume in b3

30 Magnesium Alloys - Properties in Solid and Liquid States

**Figure 46.** The plot of the activation volume in b3

all temperatures in tension and compression.

deformation temperatures in tension and compression.

deformation temperatures in tension and compression.

> s\* (MPa) 20 40 60 80 100

> > AX91

s\* (MPa) 0 20 40 60 80 100

AX alloys

AX41 AX91

s\* (MPa) 0 20 40 60 80 100

**Figure 47.** The plot of the activation volume in b3 against the effective stress σ\* estimated for AX41 and AX91 alloys at

where p and q are phenomenological parameters reflecting the shape of a resistance obstacle profile. The possible ranges of values p and q are limited by the conditions 0 < p≤ 1 and 1 ≤

C 25°C C 50°C C 100°C C 150°C C 200°C C 250°C C 300°C T 100°C T 150°C T 200°C

against the effective stress σ\* estimated for AX41 alloy at various

against the effective stress σ\* estimated for AX91 alloy at various

C 25 °C C 50 °C C100 °C C150 °C C200 °C C250 °C C300 °C T150 °C T200 °C T250 °C

$$V = \frac{\Delta G\_0 pq}{\sigma\_0^\*} \left[ 1 - \left( \frac{\sigma^\*}{\sigma\_0^\*} \right)^p \right]^{q-1} \left( \frac{\sigma^\*}{\sigma\_0^\*} \right)^{p-1} \,. \tag{14}$$

The values of the activation volume should lie at the curve given by the equation (14). Results showing all values of the activation volumes being the same for alloys of AJ as well as AX series indicate that the thermal activation is not affected by various concentrations of solute atoms. It may be concluded that the thermally activated process(-es) is(are) determined with the dislocation motion and the solute atom role is less important. The values of the activation volume may help to identify thermally activated processes considering some of the common short−range barriers to dislocation motion [27]. The dislocation–dislocation interaction mechanism has an activation volume ranging from about 102 –104 b3 , with the activation volume and enthalpy varying with strain. Couret and Caillard [28, 29] studied prismatic slip in magnesium in a wide temperature range using in situ experiments in TEM. They have reported that screw dislocations with the Burgers vector 1 / 3 112 ¯0 are able to glide on prismatic planes and their mobility is much lower than the mobility of edge dislocations. The deformation is controlled by thermally activated glide of those screw dislocation segments. A single control‐ ling mechanism was identified as the Friedel−Escaig cross slip mechanism. This mechanism assumes dissociated dislocations on compact planes, like (0001), that joint together along a critical length *L*r producing double kinks on non-compact planes. The activation volume is proportional to the critical length between two kinks. Amadieh et al. [30] found for the activation volume of the Friedel−Escaig mechanism a value of 70 b3 . Prismatic slip was also observed by Koike and Ohyama [31] in deformed AZ61 sheets. The activation of prismatic slip and subsequent annihilation of dislocation segments with opposite sign are probably the main reason for the observed internal stress decrease. The double cross slip may be a thermally activated process controlling the dislocation velocity. Beside this mechanism, the thermally activated glide of <*c+a*> dislocations should be taken into account. Mathis et al. [32] investi‐ gated the evolution of non−basal dislocations as a function of temperature in magnesium by X−ray diffraction. They found a majority of <*a*> dislocations in the as−cast state. During plastic deformation in tension the <*a*>−type dislocations remain dominant, however, the dislocation density increased by about a factor of three up to about 100 °C. At higher temperatures the fraction of <**c**+*a*>−type dislocations increased at the cost of <*a*>−type dislocations and the increase of the dislocation density is strongly reduced. The internal stress acting on dislocations is determined by the details of the internal structure at that moment and it is independent of the applied stress. The stress that changes when the applied stress is changed is only the effective stress. The internal stresses during plastic deformation of the alloys investigated here can be considered as the sum of stresses resulting from various dislocation arrangements and obstacles existing in the deformed material [33, 34]. At higher temperatures the solute atoms may diffuse to stacking fault and may influence double cross slip from basal to non-basal planes.

## **7. Dynamic strain ageing (DSA) in magnesium alloys**

Plastic deformation of alloys exhibits many phenomena associated with solute strengthening. When solute atoms can move (they may diffuse) during plastic deformation the microstructure of the deformed alloy is unstable. This microstructure instability is due to solute atoms diffusion towards to moving dislocations– the dynamic strain ageing (DSA) effect. The segregation of solute atoms at dislocations results in many phenomena:


Magnesium alloys exhibit dynamic strain ageing effects at relatively low temperatures. Portevin-Le Châtelier effect (PLC) was observed during plastic deformation of some Mg alloys [35-39]. The unstable microstructure of an alloy can influence the deformation behaviour of the alloy. It should be mentioned that room temperature is high enough to help invoking strain ageing processes in magnesium alloys.

#### **7.1. Portevin-Le Châtelier effect**

The Portevin-Le Châtelier (PLC) effect is a consequence of a complicated nature of the dislocation dynamics in metals, which depends on many structural parameters as the type of structure, grain size, texture, concentration and distribution of solute atoms. Plastic deforma‐ tion occurs inhomogeneously on the microscopic scale due to thermally activated dislocation motion through a field of obstacles. Deformation inhomogeneities in time and space, observed experimentally, are caused by collective dislocation motion. These phenomena of unstable plastic deformation are associated with the sharp localised deformation bands. Three types of PLC bands have been found. The continuously propagating type A, intermittently propagating type B with regular stress drops and stochastically nucleating type C. Figure 48 shows a representative microstructure of the undeformed Mg+0.7wt.%Nd sample after T6 temper (homogenisation at 525 °C for 5h, then quenching into water 65 °C warm with subsequent precipitation treating for 8 h at 204 °C).

**Figure 48.** Microstructure of Mg-Nd alloy after T6 temper.

increase of the dislocation density is strongly reduced. The internal stress acting on dislocations is determined by the details of the internal structure at that moment and it is independent of the applied stress. The stress that changes when the applied stress is changed is only the effective stress. The internal stresses during plastic deformation of the alloys investigated here can be considered as the sum of stresses resulting from various dislocation arrangements and obstacles existing in the deformed material [33, 34]. At higher temperatures the solute atoms may diffuse to stacking fault and may influence double cross slip from basal to non-basal

Plastic deformation of alloys exhibits many phenomena associated with solute strengthening. When solute atoms can move (they may diffuse) during plastic deformation the microstructure of the deformed alloy is unstable. This microstructure instability is due to solute atoms diffusion towards to moving dislocations– the dynamic strain ageing (DSA) effect. The

**•** local maximum in the temperature dependence of the activation volume or stress sensitivity

Magnesium alloys exhibit dynamic strain ageing effects at relatively low temperatures. Portevin-Le Châtelier effect (PLC) was observed during plastic deformation of some Mg alloys [35-39]. The unstable microstructure of an alloy can influence the deformation behaviour of the alloy. It should be mentioned that room temperature is high enough to help invoking strain

The Portevin-Le Châtelier (PLC) effect is a consequence of a complicated nature of the dislocation dynamics in metals, which depends on many structural parameters as the type of structure, grain size, texture, concentration and distribution of solute atoms. Plastic deforma‐ tion occurs inhomogeneously on the microscopic scale due to thermally activated dislocation motion through a field of obstacles. Deformation inhomogeneities in time and space, observed experimentally, are caused by collective dislocation motion. These phenomena of unstable plastic deformation are associated with the sharp localised deformation bands. Three types of PLC bands have been found. The continuously propagating type A, intermittently propagating type B with regular stress drops and stochastically nucleating type C. Figure 48 shows a representative microstructure of the undeformed Mg+0.7wt.%Nd sample after T6 temper

**7. Dynamic strain ageing (DSA) in magnesium alloys**

segregation of solute atoms at dislocations results in many phenomena:

**•** negative strain rate sensitivity in a certain temperature range,

**•** positive or non-monotonous dependence of the flow stress on temperature,

planes.

**•** post relaxation effect,

**•** the Portevin-Le Châtelier effect.

32 Magnesium Alloys - Properties in Solid and Liquid States

ageing processes in magnesium alloys.

**7.1. Portevin-Le Châtelier effect**

parameter,

The true stress-true strain curves obtained at various temperatures are presented in Fig. 49 for tension and in Fig. 50 for compression tests. It can be seen a discontinuous character of curves obtained at elevated temperatures. Serrated yielding was observed at temperatures from 200 °C to 300 °C in tension and from 200 °C to 250 °C in compression. In compression tests the deformation at 300 °C was already smooth. While the shape of the serrations obtained in tension has character A, serrations found in compression have another shape, which is more of type B. PLC effect was also observed in an AZ91 (Mg-9Al-1Zn) alloy after thermal treatment T4 (homogenisation at 413 °C for 18 h, then quenching into water of ambient temperature). Tensile and compression tests were performed over a wide temperature range from 14 to 100 °C at an initial strain rate ranged in the 10-4 s-1. The temperature dependence of the yield stress of Mg+0.7Nd exhibits a local maximum – a stress-hump (Fig. 51).

**Figure 49.** True stress-true strain curves of Mg+0.7Nd alloy deformed in tension.

**200μm**

temper (homogenisation at 525 °C for 5h, then quenching into water 65 °C warm with subsequent precipitation treating for 8 h at 204

T4 (homogenisation at 413 °C for 18 h, then quenching into water of ambient temperature). Tensile and compression tests were

 0.00 0.04 0.08 0.12 0.16 0.20

Fig. 50: True stress–true strain curves of

RT

250°C

Mg+0.7Nd

200°C

dependence of the yield stress of Mg+0.7Nd exhibits a local maximum – a stress-hump (Fig. 51).

(MPa)

various temperatures.

 0.00 0.02 0.04 0.06 0.08 0.10 0.12

Fig. 49: True stress-true strain curves of Mg+0.7Nd alloy deformed in tension.

RT 200°C

100°C

elongation of the sample during the compression test.

250°C

300°C

°C).

(MPa)

0

20

40

60

80

100

Mg+0.7Nd

120

**Figure 50.** True stress–true strain curves of Mg+0.7 alloy deformed in compression at various temperatures. performed over a wide temperature range from 14 to 100 °C at an initial strain rate ranged in the 10-4 s-1. The temperature

Mg+0.7 alloy deformed in compression at yield stress of Mg+0.7Nd alloy deformed **Figure 51.** Temperature dependence of the yield stress of Mg+0.7Nd alloy deformed in tension and compression.

in tension and compression.

Fig. 51: Temperature dependence of the

The stress-strain curve of AZ91 alloy obtained in tension at 292 K is given in Fig. 52 together with the strain dependence of the strain hardening rate = d/d. Serrations on stress-strain curves were observed at temperatures from room temperature up to 100 °C, the stress-strain curve obtained at temperature 150 °C exhibited no serrations; it was smooth. Similar experiments were performed in compression at an approximately same strain rate at temperatures from 15 to 100 °C. Character of serrations in compression is different in comparison with the tension tests as it is obvious from Fig.53. The stepwise character of curves indicates sudden The stress-strain curve of AZ91 alloy obtained in tension at 21 °C is given in Fig. 52 together with the strain dependence of the strain hardening rateθ=dσ/dε. Serrations on stress-strain curves were observed at temperatures from room temperature up to 100 °C, the stress-strain curve obtained at temperature 150 °C exhibited no serrations; it was smooth. Similar experi‐ ments were performed in compression at an approximately same strain rate at temperatures from 15 to 100 °C. Character of serrations in compression is different in comparison with the tension tests as it is obvious from Fig.53. The stepwise character of curves indicates sudden elongation of the sample during the compression test.

Influence of Solute Atoms on Deformation Behaviour of Selected Magnesium Alloys http://dx.doi.org/10.5772/58949 35

**Figure 52.** Stress-strain curve in tension obtained at 21 °C together with the strain dependence of the strain hardening rate.

**Figure 50.** True stress–true strain curves of Mg+0.7 alloy deformed in compression at various temperatures.

Mg+0.7Nd

in tension and compression.

Temperature (°C) 0 50 100 150 200 250 300 350

Fig. 51: Temperature dependence of the yield stress of Mg+0.7Nd alloy deformed

The stress-strain curve of AZ91 alloy obtained in tension at 21 °C is given in Fig. 52 together with the strain dependence of the strain hardening rateθ=dσ/dε. Serrations on stress-strain curves were observed at temperatures from room temperature up to 100 °C, the stress-strain curve obtained at temperature 150 °C exhibited no serrations; it was smooth. Similar experi‐ ments were performed in compression at an approximately same strain rate at temperatures from 15 to 100 °C. Character of serrations in compression is different in comparison with the tension tests as it is obvious from Fig.53. The stepwise character of curves indicates sudden

**Figure 51.** Temperature dependence of the yield stress of Mg+0.7Nd alloy deformed in tension and compression.

compression tension

temper (homogenisation at 525 °C for 5h, then quenching into water 65 °C warm with subsequent precipitation treating for 8 h at 204

The true stress-true strain curves obtained at various temperatures are presented in Fig. 49 for tension and in Fig. 50 for compression tests. It can be seen a discontinuous character of curves obtained at elevated temperatures. Serrated yielding was observed at temperatures from 200 °C to 300 °C in tension and from 200 °C to 250 °C in compression. In compression tests the deformation at 300 °C was already smooth. While the shape of the serrations obtained in tension has character A, serrations found in compression have another shape, which is more of type B. PLC effect was also observed in an AZ91 (Mg-9Al-1Zn) alloy after thermal treatment T4 (homogenisation at 413 °C for 18 h, then quenching into water of ambient temperature). Tensile and compression tests were performed over a wide temperature range from 14 to 100 °C at an initial strain rate ranged in the 10-4 s-1. The temperature

> 0.00 0.04 0.08 0.12 0.16 0.20

 Fig. 50: True stress–true strain curves of Mg+0.7 alloy deformed in compression at

The stress-strain curve of AZ91 alloy obtained in tension at 292 K is given in Fig. 52 together with the strain dependence of the strain hardening rate = d/d. Serrations on stress-strain curves were observed at temperatures from room temperature up to 100 °C, the stress-strain curve obtained at temperature 150 °C exhibited no serrations; it was smooth. Similar experiments were performed in compression at an approximately same strain rate at temperatures from 15 to 100 °C. Character of serrations in compression is different in comparison with the tension tests as it is obvious from Fig.53. The stepwise character of curves indicates sudden

RT

250°C

Mg+0.7Nd

200°C

**200μm**

34 Magnesium Alloys - Properties in Solid and Liquid States

02 (MPa)

elongation of the sample during the compression test.

Fig. 48: Microstructure of Mg-Nd alloy after T6 temper.

dependence of the yield stress of Mg+0.7Nd exhibits a local maximum – a stress-hump (Fig. 51).

(MPa)

various temperatures.

 0.00 0.02 0.04 0.06 0.08 0.10 0.12

Fig. 49: True stress-true strain curves of Mg+0.7Nd alloy deformed in tension.

RT 200°C

100°C

elongation of the sample during the compression test.

250°C

300°C

°C).

(MPa)

0

20

40

60

80

100

Mg+0.7Nd

120

**Figure 53.** Stress strain curve and strain dependence of the strain hardening rate obtained in compression at 100 °C.

Lavrentev [40] have reported an important role of the dislocation forest in *hcp* structures with the main basal slip. Forest dislocations, which density was increased by prior movement of dislocation avalanche, are local obstacles for moving dislocations. Cores of dislocations waiting at forest dislocations during thermally activated motion in the slip plane may be occupied by solutes, which are movable due to pipe diffusion. These smaller dislocation groups of waiting dislocations are probably spread over a set of adjacent slip planes, rather than being strictly coplanar. They produce local stress concentrations, which may be a release pulse for break‐ away of dislocation pile ups. Precursor behaviour of the local strain associated with local stress relaxation prior to the strain avalanche may be characterised by the stress drop. This type of unstable flow is characteristic of the PLC flow, in which a drop in the load is possible if the strain rate during the formation of the slip bands exceeds the strain rate imposed by the tensile testing machine. A high level of the internal stress allow pass of these deformation bands through the whole sample. As each PLC band runs through the gauge length, the formation of new forest dislocations ensures that the process repeats itself. As the dislocation movement in the slip bands ceases or becomes too slow for the applied strain rate, for instance after a drop in the load or at the end of a PLC band, the applied stress must increase in order to resume deformation. Activation of new sources can only occurs by cutting through the forest dislo‐ cations created by secondary slip around slip bands no longer active. The forest dislocations increase the critical resolved shear stress for dislocation motion, but provide little strain hardening. If the activation of a source occurs under conditions of stable planar glide, dislo‐ cations can form a dynamic pile up able to move at increasing speed and decreasing levels of the applied stress because of the development of stress concentration ahead of the moving dislocations. This may spontaneously lead to an avalanche of dislocations that meets the general requirement for unstable tensile flow, since a drop in the load, as observed during the PLC effect, is possible only if the strain rate during the formation of the slip bands exceeds the strain rate imposed by the tensile testing machine. An additional factor is the following: as dislocations pile up in front of obstacles, the average dislocation velocity in the glide plane becomes very low so that moving solute atoms may diffuse towards the dislocations and pin them down.

#### **7.2. Other manifestations of the dynamic strain ageing**

The temperature dependence of the yield stress, σ02, of Mg+0.7Nd alloy deformed in tension and compression (Fig. 51) shows that the course of the temperature dependence is nonmonotonous. A local maximum in the temperature dependence of the yield stress obtained in compression and tension is observed in a temperature range of 150 – 250 °C.

Similar local minimum has been observed in the temperature dependence of the yield stress, σ02, and the maximum compressive stress, σmax, of ZE41 (Mg-4Zn-1RE) magnesium alloy as it can be seem in Fig. 54. While the local maximum in the temperature dependence is observed in alloys containing rare earths in the temperature interval 150-250 °C, alloys containing Al exhibit this local maximum between room temperature and 100 °C. This is demonstrated in Figs. 55 and 56 for AZ63 (Mg-6Al-3Zn) and AJ51 (Mg-5Al-0.6Sr), respectively. The strain rate dependence of the yield stress of Mg+0.7Nd alloy deformed in tension at 250 °C is presented Influence of Solute Atoms on Deformation Behaviour of Selected Magnesium Alloys http://dx.doi.org/10.5772/58949 37

σ02 (MPa)

AZ63

σ02 (MPa)

AJ51

temperature (°C) 0 50 100 150 200 250 300 350

Fig. 56: Temperature dependence of the

yield stress of AJ51 alloy.

tension

293 K

ε 0.10 0.11 0.12 0.13

Fig. 59: Sequences of the stress strain curves of AZ91 obtained for two strain rates. Negative strain rate sensitivity is

3.6x10-5s-1

1.8x10-6s-1

AZ91 T4

temperature (°C) 0 50 100 150 200 250 300 350

Fig. 55: Temperature dependence of the

the yield stress is practically independent of the strain rate. Again, the results indicate some dynamic strain ageing. It is also obvious

the stress strain curves of AZ91 alloy deformed at room temperature and two various strain rates are shown in Fig.59. It should be

ZE41

 50°C 150°C 100°C

σ (MPa)

280

obvious.

290

300

310

320

σ02 [MPa]

60

three temperatures:

80

100

120

140

160

dε/dt [s-1] 10-5 10-4 10-3 10-2 10-1

Very effective tool for studying of strain ageing phenomena are the stress relaxation tests. The SR curves are usually analysed assuming that the mobile dislocation density ρm and internal stress σi are constant during the SR test. An unstable structure, changes in the mobile dislocation density and/or in the internal stress, may influence the course of the SR. In some SR tests, we observed a post relaxation effect. The flow stress at the beginning of deformation after stress relaxation σ1 was higher (or lower) than the stress at the beginning of the stress relaxation σ(0). The values of ∆σ = σ(0) - σ1 depend on strain, starting stress of the SR test and temperature. Examples of such dependences can be seen in Figs. 60-63 for magnesium alloys containing rare earth elements – Mg+0.7Nd alloy, ZE41 alloy, AE42 (Mg-4Al-2RE) and QE22 (Mg-2Ag-2RE). Similar dependences for alloys of the AZ series are presented in Figs. 64 and 65. It is obvious that the post relaxation effect is limited only on certain stress (strain) and temperature ranges. Local maxima are observed in these dependences. Such experimental results indicate that these phenomena are related to the

Fig. 58: Strain rate dependence of the yield stress of ZE41 alloy deformed at

yield stress of AZ63 alloy.

 Fig. 54: Temperature dependence of the **Figure 54.** Temperature dependence of the yield and maximum stresses of ZE41 alloy

Lavrentev [40] have reported an important role of the dislocation forest in *hcp* structures with the main basal slip. Forest dislocations, which density was increased by prior movement of dislocation avalanche, are local obstacles for moving dislocations. Cores of dislocations waiting at forest dislocations during thermally activated motion in the slip plane may be occupied by solutes, which are movable due to pipe diffusion. These smaller dislocation groups of waiting dislocations are probably spread over a set of adjacent slip planes, rather than being strictly coplanar. They produce local stress concentrations, which may be a release pulse for break‐ away of dislocation pile ups. Precursor behaviour of the local strain associated with local stress relaxation prior to the strain avalanche may be characterised by the stress drop. This type of unstable flow is characteristic of the PLC flow, in which a drop in the load is possible if the strain rate during the formation of the slip bands exceeds the strain rate imposed by the tensile testing machine. A high level of the internal stress allow pass of these deformation bands through the whole sample. As each PLC band runs through the gauge length, the formation of new forest dislocations ensures that the process repeats itself. As the dislocation movement in the slip bands ceases or becomes too slow for the applied strain rate, for instance after a drop in the load or at the end of a PLC band, the applied stress must increase in order to resume deformation. Activation of new sources can only occurs by cutting through the forest dislo‐ cations created by secondary slip around slip bands no longer active. The forest dislocations increase the critical resolved shear stress for dislocation motion, but provide little strain hardening. If the activation of a source occurs under conditions of stable planar glide, dislo‐ cations can form a dynamic pile up able to move at increasing speed and decreasing levels of the applied stress because of the development of stress concentration ahead of the moving dislocations. This may spontaneously lead to an avalanche of dislocations that meets the general requirement for unstable tensile flow, since a drop in the load, as observed during the PLC effect, is possible only if the strain rate during the formation of the slip bands exceeds the strain rate imposed by the tensile testing machine. An additional factor is the following: as dislocations pile up in front of obstacles, the average dislocation velocity in the glide plane becomes very low so that moving solute atoms may diffuse towards the dislocations and pin

36 Magnesium Alloys - Properties in Solid and Liquid States

them down.

**7.2. Other manifestations of the dynamic strain ageing**

The temperature dependence of the yield stress, σ02, of Mg+0.7Nd alloy deformed in tension and compression (Fig. 51) shows that the course of the temperature dependence is nonmonotonous. A local maximum in the temperature dependence of the yield stress obtained in

Similar local minimum has been observed in the temperature dependence of the yield stress, σ02, and the maximum compressive stress, σmax, of ZE41 (Mg-4Zn-1RE) magnesium alloy as it can be seem in Fig. 54. While the local maximum in the temperature dependence is observed in alloys containing rare earths in the temperature interval 150-250 °C, alloys containing Al exhibit this local maximum between room temperature and 100 °C. This is demonstrated in Figs. 55 and 56 for AZ63 (Mg-6Al-3Zn) and AJ51 (Mg-5Al-0.6Sr), respectively. The strain rate dependence of the yield stress of Mg+0.7Nd alloy deformed in tension at 250 °C is presented

compression and tension is observed in a temperature range of 150 – 250 °C.

yield and maximum stresses of ZE41 alloy

that the values of the yield stress at 100 ºC are lower than those at 150 ºC at all imposed strain rates. Examples of short sequences of **Figure 55.** Temperature dependence of the yield stress of AZ63 alloy.

strain rate dε/dt s-1)

 Fig. 57: Strain rate dependence of the yield stress of Mg+0.7Nd alloy at two

dislocations and solute atoms mobility.

10-4 10-3 20 **Figure 56.** Temperature dependence of the yield stress of AJ51 alloy.

temperatures.

ZE41

Mg+0.7Nd

strain rate dε/dt s-1) 10-4 10-3

yield stress of Mg+0.7Nd alloy at two

dislocations and solute atoms mobility.

250 °C 300 °C

300 σ<sup>02</sup>

σ (MPa)

yield stress (MPa)

temperatures.

0

100

200

σ (MPa)

0

100

100

200

in Fig.57. Negative strain rate sensitivity was found for three strain rates from 5.5x10-5 to 5.5x10-4 s-1. The strain rate dependence of the yield stress of ZE41 alloy deformed at three temperatures is given in Fig. 58. It can be seen that the strain rate dependences have an anomalous course. The yield stress usually increases with increasing strain rate. In this case, the yield stress increases only at lower strain rates. At strain rates higher than 10-4 s-1, the yield stress decreases with increasing strain rate at 50 and 150 ºC. At 100 ºC the yield stress is practically independent of the strain rate. Again, the results indicate some dynamic strain ageing. It is also obvious that the values of the yield stress at 100 ºC are lower than those at 150 ºC at all imposed strain rates. Examples of short sequences of the stress strain curves of AZ91 alloy deformed at room temperature and two various strain rates are shown in Fig.59. It should be mentioned that the curves are not shifted; accordingly the negative strain rate sensitivity is obvious. dependence is observed in alloys containing rare earths in the temperature interval 150-250 °C, alloys containing Al exhibit this local maximum between room temperature and 100 °C. This is demonstrated in Figs. 55 and 56 for AZ63 (Mg-6Al-3Zn) and AJ51 (Mg-5Al-0.6Sr), respectively. The strain rate dependence of the yield stress of Mg+0.7Nd alloy deformed in tension at 250 °C is presented in Fig.57. Negative strain rate sensitivity was found for three strain rates from 5.5x10-5 to 5.5x10-4 s-1. The strain rate dependence of the yield stress of ZE41 alloy deformed at three temperatures is given in Fig. 58. It can be seen that the strain rate dependences have an anomalous course. The yield stress usually increases with increasing strain rate. In this case, the yield stress increases only at lower strain rates. At strain rates higher than 10-4 s-1, the yield stress decreases with increasing strain rate at 50 and 150 °C. At 100 °C the yield stress is practically independent of the strain rate. Again, the results indicate some dynamic strain ageing. It is also obvious that the values of the yield stress at 100 ºC are lower than those at 150 ºC at all imposed strain rates. Examples of short sequences of the stress strain curves of AZ91 alloy deformed at room temperature and two various strain rates are shown in Fig.59. It should be mentioned that the curves are not shifted; accordingly the negative strain rate sensitivity is obvious. temperature (°C) 0 50 100 150 200 250 300 350 σmax Fig. 54: Temperature dependence of the yield and maximum stresses of ZE41 alloy temperature (°C) 0 50 100 150 200 250 300 350 σ02 (MPa) 40 50 60 70 80 90 Fig. 55: Temperature dependence of the yield stress of AZ63 alloy. 0 50 100 150 200 250 300 350 σ02 (MPa) 20 40 60 80 100 120 Fig. 56: Temperature dependence of the yield stress of AJ51 alloy.

ZE41

σmax

AZ63

σ02 (MPa)

140 160 AZ63

σ02 (MPa)

AJ51

temperature (°C) 0 50 100 150 200 250 300 350

Fig. 56: Temperature dependence of the

yield stress of AJ51 alloy.

tension

293 K

ε 0.10 0.11 0.12 0.13

Fig. 59: Sequences of the stress strain curves of AZ91 obtained for two strain rates. Negative strain rate sensitivity is

3.6x10-5s-1

1.8x10-6s-1

AZ91 T4

temperature (°C) 0 50 100 150 200 250 300 350

AJ51

Fig. 55: Temperature dependence of the

stress, σmax, of ZE41 (Mg-4Zn-1RE) magnesium alloy as it can be seem in Fig. 54. While the local maximum in the temperature

temperature (°C)

ZE41

 50°C 150°C 100°C

σ (MPa)

280

obvious.

290

300

310

320

σ02 [MPa]

σ (MPa)

280

obvious.

290

300

310

320

60

three temperatures:

tension

293 K

80

100

120

140

160

dε/dt [s-1] 10-5 10-4 10-3 10-2 10-1

3.6x10-5s-1

1.8x10-6s-1

AZ91 T4

Mg+0.7Nd alloy, ZE41 alloy, AE42 (Mg-4Al-2RE) and QE22 (Mg-2Ag-2RE). Similar dependences for alloys of the AZ series are

ε 0.10 0.11 0.12 0.13

ranges. Local maxima are observed in these dependences. Such experimental results indicate that these phenomena are related to the

curves of AZ91 obtained for two strain rates. Negative strain rate sensitivity is

Fig. 58: Strain rate dependence of the yield stress of ZE41 alloy deformed at

yield stress of AZ63 alloy.

300 σ<sup>02</sup>

temperature (°C) 0 50 100 150 200 250 300 350

 Fig. 54: Temperature dependence of the yield and maximum stresses of ZE41 alloy

 Fig. 57: Strain rate dependence of the **Figure 57.** Strain rate dependence of the yield stress of Mg+0.7Nd alloy at two temperatures. mentioned that the curves are not shifted; accordingly the negative strain rate sensitivity is obvious.

dislocations and solute atoms mobility.

Very effective tool for studying of strain ageing phenomena are the stress relaxation tests. The SR curves are usually analysed assuming that the mobile dislocation density ρm and internal stress σi are constant during the SR test. An unstable structure, changes in the mobile dislocation density and/or in the internal stress, may influence the course of the SR. In some SR tests, we observed a post relaxation effect. The flow stress at the beginning of deformation after stress relaxation σ1 was higher (or lower) than the stress at the beginning of the stress relaxation σ(0). The values of ∆σ = σ(0) - σ1 depend on strain, starting stress of the SR test and temperature. Examples of such dependences can be seen in Figs. 60-63 for magnesium alloys containing rare earth elements – Mg+0.7Nd alloy, ZE41 alloy, AE42 (Mg-4Al-2RE) and QE22 (Mg-2Ag-2RE). Similar dependences for alloys of the AZ series are presented in Figs. 64 and 65. It is obvious that the post relaxation effect is limited only on certain stress (strain) and temperature ranges. Local maxima are observed in these dependences. Such experimental results indicate that these phenomena are related to the

yield stress of ZE41 alloy deformed at

yield stress of Mg+0.7Nd alloy at two

presented in Figs. 64 and 65. It is obvious that the post relaxation effect is limited only on certain stress (strain) and temperature Fig. 57: Strain rate dependence of the Fig. 58: Strain rate dependence of the Fig. 59: Sequences of the stress strain **Figure 58.** Strain rate dependence of the yield stress of ZE41 alloy deformed at three temperatures:

three temperatures:

AJ51

temperature (°C) 0 50 100 150 200 250 300 350

Fig. 56: Temperature dependence of the

yield stress of AJ51 alloy.

in Fig.57. Negative strain rate sensitivity was found for three strain rates from 5.5x10-5 to 5.5x10-4 s-1. The strain rate dependence of the yield stress of ZE41 alloy deformed at three temperatures is given in Fig. 58. It can be seen that the strain rate dependences have an anomalous course. The yield stress usually increases with increasing strain rate. In this case, the yield stress increases only at lower strain rates. At strain rates higher than 10-4 s-1, the yield stress decreases with increasing strain rate at 50 and 150 ºC. At 100 ºC the yield stress is practically independent of the strain rate. Again, the results indicate some dynamic strain ageing. It is also obvious that the values of the yield stress at 100 ºC are lower than those at 150 ºC at all imposed strain rates. Examples of short sequences of the stress strain curves of AZ91 alloy deformed at room temperature and two various strain rates are shown in Fig.59. It should be mentioned that the curves are not shifted; accordingly the negative strain rate sensitivity is

Mg+0.7Nd

temperature (°C) 0 50 100 150 200 250 300 350

Fig. 55: Temperature dependence of the

Similar local minimum has been observed in the temperature dependence of the yield stress, σ02, and the maximum compressive

maximum between room temperature and 100 °C. This is demonstrated in Figs. 55 and 56 for AZ63 (Mg-6Al-3Zn) and AJ51 (Mg-

strain rate dε/dt s-1) 10-4 10-3

 Fig. 57: Strain rate dependence of the yield stress of Mg+0.7Nd alloy at two

ZE41

 50°C 150°C 100°C

dislocations and solute atoms mobility.

Very effective tool for studying of strain ageing phenomena are the stress relaxation tests. The SR curves are usually analysed assuming that the mobile dislocation density ρm and internal stress σi are constant during the SR test. An unstable structure, changes in the mobile dislocation density and/or in the internal stress, may influence the course of the SR. In some SR tests, we observed a post relaxation effect. The flow stress at the beginning of deformation after stress relaxation σ1 was higher (or lower) than the stress at the beginning of the stress relaxation σ(0). The values of ∆σ = σ(0) - σ1 depend on strain, starting stress of the SR test and temperature. Examples of such dependences can be seen in Figs. 60-63 for magnesium alloys containing rare earth elements – Mg+0.7Nd alloy, ZE41 alloy, AE42 (Mg-4Al-2RE) and QE22 (Mg-2Ag-2RE). Similar dependences for alloys of the AZ series are presented in Figs. 64 and 65. It is obvious that the post relaxation effect is limited only on certain stress (strain) and temperature ranges. Local maxima are observed in these dependences. Such experimental results indicate that these phenomena are related to the

Fig. 58: Strain rate dependence of the yield stress of ZE41 alloy deformed at

**Figure 58.** Strain rate dependence of the yield stress of ZE41 alloy deformed at three temperatures:

dε/dt [s-1] 10-5 10-4 10-3 10-2 10-1

250 °C 300 °C

yield stress of AZ63 alloy.

ZE41

σmax

AZ63

σ02 (MPa)

mentioned that the curves are not shifted; accordingly the negative strain rate sensitivity is obvious.

σ02 (MPa)

σ02 [MPa]

σ (MPa)

280

obvious.

290

300

310

320

60

three temperatures:

tension

293 K

80

100

120

140

160

AZ63

σ02 (MPa)

temperature (°C) 0 50 100 150 200 250 300 350

AJ51

Fig. 55: Temperature dependence of the

stress, σmax, of ZE41 (Mg-4Zn-1RE) magnesium alloy as it can be seem in Fig. 54. While the local maximum in the temperature

temperature (°C) 0 50 100 150 200 250 300 350

Fig. 56: Temperature dependence of the

yield stress of AJ51 alloy.

ZE41

 50°C 150°C 100°C

obvious.

dε/dt [s-1]

3.6x10-5s-1

1.8x10-6s-1

AZ91 T4

ranges. Local maxima are observed in these dependences. Such experimental results indicate that these phenomena are related to the

curves of AZ91 obtained for two strain rates. Negative strain rate sensitivity is

ε 0.10 0.11 0.12 0.13

Fig. 58: Strain rate dependence of the yield stress of ZE41 alloy deformed at

yield stress of AZ63 alloy.

300 σ<sup>02</sup>

σ (MPa)

yield stress (MPa)

mentioned that the curves are not shifted; accordingly the negative strain rate sensitivity is obvious.

temperatures.

60

three temperatures:

80

100

120

140

160

σ02 [MPa]

**Figure 57.** Strain rate dependence of the yield stress of Mg+0.7Nd alloy at two temperatures.

0

σ02 (MPa)

100

200

temperature (°C) 0 50 100 150 200 250 300 350

 Fig. 54: Temperature dependence of the yield and maximum stresses of ZE41 alloy

obvious.

Mg+0.7Nd

strain rate dε/dt s-1) 10-4 10-3

 Fig. 57: Strain rate dependence of the yield stress of Mg+0.7Nd alloy at two

dislocations and solute atoms mobility.

250 °C 300 °C

ZE41

σmax

38 Magnesium Alloys - Properties in Solid and Liquid States

300 σ<sup>02</sup>

σ (MPa)

yield stress (MPa)

temperatures.

0

100

200

temperature (°C) 0 50 100 150 200 250 300 350

 Fig. 54: Temperature dependence of the yield and maximum stresses of ZE41 alloy

**Figure 59.** Sequences of the stress strain curves of AZ91 obtained for two strain rates. Negative strain rate sensitivity is obvious.

10-5 10-4 10-3 10-2 10-1 tension ε 0.10 0.11 0.12 0.13 σ (MPa) 280 290 300 310 320 3.6x10-5s-1 1.8x10-6s-1 293 K AZ91 T4 Fig. 59: Sequences of the stress strain stress, σmax, of ZE41 (Mg-4Zn-1RE) magnesium alloy as it can be seem in Fig. 54. While the local maximum in the temperature dependence is observed in alloys containing rare earths in the temperature interval 150-250 °C, alloys containing Al exhibit this local 5Al-0.6Sr), respectively. The strain rate dependence of the yield stress of Mg+0.7Nd alloy deformed in tension at 250 °C is presented in Fig.57. Negative strain rate sensitivity was found for three strain rates from 5.5x10-5 to 5.5x10-4 s-1. The strain rate dependence of the yield stress of ZE41 alloy deformed at three temperatures is given in Fig. 58. It can be seen that the strain rate dependences have an anomalous course. The yield stress usually increases with increasing strain rate. In this case, the yield stress increases only at lower strain rates. At strain rates higher than 10-4 s-1, the yield stress decreases with increasing strain rate at 50 and 150 °C. At 100 °C the yield stress is practically independent of the strain rate. Again, the results indicate some dynamic strain ageing. It is also obvious that the values of the yield stress at 100 ºC are lower than those at 150 ºC at all imposed strain rates. Examples of short sequences of the stress strain curves of AZ91 alloy deformed at room temperature and two various strain rates are shown in Fig.59. It should be Very effective tool for studying of strain ageing phenomena are the stress relaxation tests. The SR curves are usually analysed assuming that the mobile dislocation density ρm and internal stress σ<sup>i</sup> are constant during the SR test. An unstable structure, changes in the mobile disloca‐ tion density and/or in the internal stress, may influence the course of the SR. In some SR tests, we observed a post relaxation effect. The flow stress at the beginning of deformation after stress relaxation *σ*1 was higher (or lower) than the stress at the beginning of the stress relaxation *σ*(0). The values of Δ*σ*=*σ*(0)-*σ*1 depend on strain, starting stress of the SR test and temperature. Examples of such dependences can be seen in Figs. 60-63 for magnesium alloys containing rare earth elements – Mg+0.7Nd alloy, ZE41 alloy, AE42 (Mg-4Al-2RE) and QE22 (Mg-2Ag-2RE). Similar dependences for alloys of the AZ series are presented in Figs. 64 and 65. It is obvious that the post relaxation effect is limited only on certain stress (strain) and temperature ranges. Local maxima are observed in these dependences. Such experimental results indicate that these phenomena are related to the dislocations and solute atoms mobility.

ZE41

150°C

AZ63

r

∆σ (MPa)

2

σ = σd + σf, (15)

]. (16)

3

4

5

6

∆σ (MPa)

0

1

2

3

4

5

AE42

temperature (°C) 80 100 120 140 160 180 200 220

temperature (°C) 0 40 80 120 160

Fig. 65: Temperature dependence of the stress increment after SR at two SR starting strains (AZ91 alloy).

ε = 0.005 ε = 0.015

Fig. 62: Temperature variation of the stress increment after SR at two starting

strains of SR (AE42 alloy).

0.005 0.01 0.03

AZ91 T6

σ (MPa) 100 150 200 250

SR at two temperatures (ZE41 alloy).

Solute atoms become mobile with increasing temperature. During a stress relaxation test, the dislocation velocity decreases, and at higher temperatures, the mobility of foreign atoms may be close to that of the dislocations. Thus, atmospheres of foreign atoms on dislocations may form. The dislocations are pinned by the solutes and hence, in order to restart their motion, they must be freed from the atmospheres after stress relaxation. Macroscopically, this results in a yield point due to dynamic strain ageing. The stress at the beginning deformation after the SR is higher than the stress at the beginning of the stress relaxation, which is observed

> ε 0.00 0.03 0.06 0.09 0.12 0.15

Fig. 64: Strain dependence of the stress increment after SR at two temperatures

whereσ d is the dislocation component due to strong obstacles (e.g. forest dislocations, precipitates, grain boundaries) and σf is the friction stress due to the interaction between the solute atoms and moving dislocations. Mulford and Kocks [41] consider that the dislocation component of the flow stress rather than the friction stress is affected by dynamic strain ageing. Balík and Lukáč [42] take

∆c = c - c0 = ∆cM[1 - exp(-ta/t0)

Here c is the local solute concentration in the dislocation core, c0 is the nominal solute concentration in the matrix and ∆cM is the maximum concentration increment. The exponent r is equal 2/3 or 1/3 for bulk or pipe diffusion, respectively. The relaxation time t<sup>0</sup> depends on the binding energy between a dislocation and a solute atom, on solute concentration, and on the diffusion coefficient of solute atoms. t0 is inversely proportional to the diffusion coefficient in the case of bulk diffusion, whereas for pipe diffusion 1/t0∼

25°C 50°C

∆σ (MPa)

∆σ (MPa)

100 °C 200 °C

(AZ63 alloy).

The flow stress, necessary for the dislocation movement, may be expressed as a sum of two components

The local solute concentration increment, ∆c, on the dislocations can be expressed as [43-45]

3/2 [46], where ρf is the density of forest dislocations.

into account the influence of solute atoms on the both flow stress components in the dynamic strain ageing regime.

200°C

Mg+0.7Nd alloy, ZE41 alloy, AE42 (Mg-4Al-2RE) and QE22 (Mg-2Ag-2RE). Similar dependences for alloys of the AZ series are presented in Figs. 64 and 65. It is obvious that the post relaxation effect is limited only on certain stress (strain) and temperature Fig. 59: Sequences of the stress strain Fig. 60: Dependence of the stress increment after the SR on the starting stress Fig. 61: Dependence of the stress increment after SR on the starting stress of **Figure 60.** Dependence of the stress increment after the SR on the starting stress of SR estimated at two temperatures (Mg+0.7Nd).

<sup>14</sup> 25 °C

of SR estimated at two temperatures

Fig. 63: Strain dependence of the stress increment after SR at three temperatures

ε 0.00 0.04 0.08 0.12 0.16 0.20

(Mg+0.7Nd).

experimentally.

∆σ (MPa)

(QE22 alloy).

QE22

Dρ<sup>f</sup>

ZE41

150°C

σ (MPa) 100 150 200 250

beginning deformation after the SR is higher than the stress at the beginning of the stress relaxation, which is observed

ε 0.00 0.03 0.06 0.09 0.12 0.15

Fig. 64: Strain dependence of the stress increment after SR at two temperatures

whereσ d is the dislocation component due to strong obstacles (e.g. forest dislocations, precipitates, grain boundaries) and σf is the friction stress due to the interaction between the solute atoms and moving dislocations. Mulford and Kocks [41] consider that the dislocation component of the flow stress rather than the friction stress is affected by dynamic strain ageing. Balík and Lukáč [42] take

∆c = c - c0 = ∆cM[1 - exp(-ta/t0)

Here c is the local solute concentration in the dislocation core, c0 is the nominal solute concentration in the matrix and ∆cM is the maximum concentration increment. The exponent r is equal 2/3 or 1/3 for bulk or pipe diffusion, respectively. The relaxation time t<sup>0</sup> depends on the binding energy between a dislocation and a solute atom, on solute concentration, and on the diffusion coefficient of solute atoms. t0 is inversely proportional to the diffusion coefficient in the case of bulk diffusion, whereas for pipe diffusion 1/t0∼

25°C 50°C

AZ63

r

∆σ (MPa)

2

3

4

5

6

Mg+0.7Nd

∆σ (MPa)

∆σ (MPa)

100 °C 200 °C

0

(AZ63 alloy).

The flow stress, necessary for the dislocation movement, may be expressed as a sum of two components

The local solute concentration increment, ∆c, on the dislocations can be expressed as [43-45]

3/2 [46], where ρf is the density of forest dislocations.

into account the influence of solute atoms on the both flow stress components in the dynamic strain ageing regime.

1

2

3

4

5

6

200°C

σ (MPa) 30 40 50 60 70 80 90

increment after the SR on the starting stress

<sup>14</sup> 25 °C

150 °C 200 °C

Fig. 60: Dependence of the stress

of SR estimated at two temperatures

Fig. 63: Strain dependence of the stress increment after SR at three temperatures

ε

0.00 0.04 0.08 0.12 0.16 0.20

∆σ (MPa)


(Mg+0.7Nd).

experimentally.

∆σ (MPa)

(QE22 alloy).

QE22

Dρ<sup>f</sup>

0 1

2

3

4

5

**Figure 61.** Dependence of the stress increment after SR on the starting stress of SR at two temperatures (ZE41 alloy).

Fig. 61: Dependence of the stress increment after SR on the starting stress of Fig. 62: Temperature variation of the stress increment after SR at two starting **Figure 62.** Temperature variation of the stress increment after SR at two starting strains of SR (AE42 alloy).

SR at two temperatures (ZE41 alloy). strains of SR (AE42 alloy). Solute atoms become mobile with increasing temperature. During a stress relaxation test, the dislocation velocity decreases, and at higher temperatures, the mobility of foreign atoms may be close to that of the dislocations. Thus, atmospheres of foreign atoms on dislocations may form. The dislocations are pinned by the solutes and hence, in order to restart their motion, they must be freed from the atmospheres after stress relaxation. Macroscopically, this results in a yield point due to dynamic strain ageing. The stress at the Solute atoms become mobile with increasing temperature. During a stress relaxation test, the dislocation velocity decreases, and at higher temperatures, the mobility of foreign atoms may be close to that of the dislocations. Thus, atmospheres of foreign atoms on dislocations may form. The dislocations are pinned by the solutes and hence, in order to restart their motion, they must be freed from the atmospheres after stress relaxation. Macroscopically, this results in a yield point due to dynamic strain ageing. The stress at the beginning deformation after the SR is higher than the stress at the beginning of the stress relaxation, which is observed experimentally.

> temperature (°C) 0 40 80 120 160

Fig. 65: Temperature dependence of the stress increment after SR at two SR

starting strains (AZ91 alloy).

]. (16)

σ = σd + σf, (15)

ε = 0.005 ε = 0.015 AZ91 T6

the atmospheres after stress relaxation. Macroscopically, this results in a yield point due to dynamic strain ageing. The stress at the beginning deformation after the SR is higher than the stress at the beginning of the stress relaxation, which is observed experimentally. Influence of Solute Atoms on Deformation Behaviour of Selected Magnesium Alloys http://dx.doi.org/10.5772/58949 41

∆σ (MPa)

(AZ63 alloy).

Mg+0.7Nd

∆σ (MPa)

200°C

ZE41

150°C

AZ63

r

∆σ (MPa)

2

σ = σd + σf, (15)

]. (16)

3

4

5

6

∆σ (MPa)

0

1

2

3

4

5

AE42

temperature (°C) 80 100 120 140 160 180 200 220

temperature (°C) 0 40 80 120 160

Fig. 65: Temperature dependence of the stress increment after SR at two SR starting strains (AZ91 alloy).

ε = 0.005 ε = 0.015

Fig. 62: Temperature variation of the stress increment after SR at two starting

strains of SR (AE42 alloy).

0.005 0.01 0.03

AZ91 T6

σ (MPa) 100 150 200 250

increment after SR on the starting stress of SR at two temperatures (ZE41 alloy).

Solute atoms become mobile with increasing temperature. During a stress relaxation test, the dislocation velocity decreases, and at higher temperatures, the mobility of foreign atoms may be close to that of the dislocations. Thus, atmospheres of foreign atoms on dislocations may form. The dislocations are pinned by the solutes and hence, in order to restart their motion, they must be freed from

> ε 0.00 0.03 0.06 0.09 0.12 0.15

increment after SR at two temperatures

∆c = c - c0 = ∆cM[1 - exp(-ta/t0)

25°C 50°C

Fig. 61: Dependence of the stress

σ (MPa) 30 40 50 60 70 80 90

increment after the SR on the starting stress of SR estimated at two temperatures

150 °C 200 °C

Fig. 60: Dependence of the stress

∆σ (MPa)

(Mg+0.7Nd).

ZE41

150°C

AZ63

r

σ (MPa)

increment after SR on the starting stress of SR at two temperatures (ZE41 alloy).

> ε 0.00 0.03 0.06 0.09 0.12 0.15

increment after SR at two temperatures

whereσ d is the dislocation component due to strong obstacles (e.g. forest dislocations, precipitates, grain boundaries) and σf is the friction stress due to the interaction between the solute atoms and moving dislocations. Mulford and Kocks [41] consider that the dislocation component of the flow stress rather than the friction stress is affected by dynamic strain ageing. Balík and Lukáč [42] take

∆c = c - c0 = ∆cM[1 - exp(-ta/t0)

Here c is the local solute concentration in the dislocation core, c0 is the nominal solute concentration in the matrix and ∆cM is the maximum concentration increment. The exponent r is equal 2/3 or 1/3 for bulk or pipe diffusion, respectively. The relaxation time t<sup>0</sup> depends on the binding energy between a dislocation and a solute atom, on solute concentration, and on the diffusion coefficient of solute atoms. t0 is inversely proportional to the diffusion coefficient in the case of bulk diffusion, whereas for pipe diffusion 1/t0∼

25°C 50°C

Fig. 61: Dependence of the stress

Fig. 63: Strain dependence of the stress Fig. 64: Strain dependence of the stress **Figure 63.** Strain dependence of the stress increment after SR at three temperatures (QE22 alloy). 100 150 200 250 80 100 120 140 160 180 200 220

0

ZE41

∆σ (MPa)

(Mg+0.7Nd).

experimentally.

150°C

150 °C 200 °C

Fig. 60: Dependence of the stress

of SR estimated at two temperatures

Mg+0.7Nd

∆σ (MPa)

∆σ (MPa)

100 °C 200 °C

(AZ63 alloy).

The flow stress, necessary for the dislocation movement, may be expressed as a sum of two components

The local solute concentration increment, ∆c, on the dislocations can be expressed as [43-45]

3/2 [46], where ρf is the density of forest dislocations.

into account the influence of solute atoms on the both flow stress components in the dynamic strain ageing regime.

200°C

0.005 0.01 0.03

σ (MPa) 30 40 50 60 70 80 90

increment after the SR on the starting stress

<sup>14</sup> 25 °C

AZ91 T6

(QE22 alloy).

QE22

Fig. 63: Strain dependence of the stress increment after SR at three temperatures

ε 0.00 0.04 0.08 0.12 0.16 0.20

s (MPa) 100 150 200 250

temperature (°C) 80 100 120 140 160 180 200 220

∆σ (MPa)

temperature (°C) 0 40 80 120 160

Dρ<sup>f</sup>

Fig. 65: Temperature dependence of the stress increment after SR at two SR

starting strains (AZ91 alloy).

]. (16)

σ = σd + σf, (15)

ε = 0.005 ε = 0.015

Fig. 62: Temperature variation of the stress increment after SR at two starting

Solute atoms become mobile with increasing temperature. During a stress relaxation test, the dislocation velocity decreases, and at higher temperatures, the mobility of foreign atoms may be close to that of the dislocations. Thus, atmospheres of foreign atoms on dislocations may form. The dislocations are pinned by the solutes and hence, in order to restart their motion, they must be freed from the atmospheres after stress relaxation. Macroscopically, this results in a yield point due to dynamic strain ageing. The stress at the beginning deformation after the SR is higher than the stress at the beginning of the stress relaxation, which is observed

strains of SR (AE42 alloy).

**Figure 62.** Temperature variation of the stress increment after SR at two starting strains of SR (AE42 alloy).

**Figure 61.** Dependence of the stress increment after SR on the starting stress of SR at two temperatures (ZE41 alloy).

AE42

Ds (MPa)

40 Magnesium Alloys - Properties in Solid and Liquid States

ZE41

150°C

AZ63

experimentally.

r

∆σ (MPa)

2

3

4

5

6

σ (MPa) 100 150 200 250

increment after SR on the starting stress of SR at two temperatures (ZE41 alloy).

Solute atoms become mobile with increasing temperature. During a stress relaxation test, the dislocation velocity decreases, and at higher temperatures, the mobility of foreign atoms may be close to that of the dislocations. Thus, atmospheres of foreign atoms on dislocations may form. The dislocations are pinned by the solutes and hence, in order to restart their motion, they must be freed from the atmospheres after stress relaxation. Macroscopically, this results in a yield point due to dynamic strain ageing. The stress at the beginning deformation after the SR is higher than the stress at the beginning of the stress relaxation, which is observed

> ε 0.00 0.03 0.06 0.09 0.12 0.15

Fig. 64: Strain dependence of the stress increment after SR at two temperatures

whereσ d is the dislocation component due to strong obstacles (e.g. forest dislocations, precipitates, grain boundaries) and σf is the friction stress due to the interaction between the solute atoms and moving dislocations. Mulford and Kocks [41] consider that the dislocation component of the flow stress rather than the friction stress is affected by dynamic strain ageing. Balík and Lukáč [42] take

∆c = c - c0 = ∆cM[1 - exp(-ta/t0)

Here c is the local solute concentration in the dislocation core, c0 is the nominal solute concentration in the matrix and ∆cM is the maximum concentration increment. The exponent r is equal 2/3 or 1/3 for bulk or pipe diffusion, respectively. The relaxation time t<sup>0</sup> depends on the binding energy between a dislocation and a solute atom, on solute concentration, and on the diffusion coefficient of solute atoms. t0 is inversely proportional to the diffusion coefficient in the case of bulk diffusion, whereas for pipe diffusion 1/t0∼

25°C 50°C

Fig. 61: Dependence of the stress

Mg+0.7Nd

∆σ (MPa)

∆σ (MPa)

100 °C 200 °C

0

(AZ63 alloy).

The flow stress, necessary for the dislocation movement, may be expressed as a sum of two components

The local solute concentration increment, ∆c, on the dislocations can be expressed as [43-45]

3/2 [46], where ρf is the density of forest dislocations.

into account the influence of solute atoms on the both flow stress components in the dynamic strain ageing regime.

1

2

3

4

5

6

200°C

σ (MPa) 30 40 50 60 70 80 90

increment after the SR on the starting stress

<sup>14</sup> 25 °C

150 °C 200 °C

Fig. 60: Dependence of the stress

of SR estimated at two temperatures

Fig. 63: Strain dependence of the stress increment after SR at three temperatures

ε

0.00 0.04 0.08 0.12 0.16 0.20

∆σ (MPa)


(Mg+0.7Nd).

experimentally.

∆σ (MPa)

(QE22 alloy).

QE22

Dρ<sup>f</sup>

0 1

2

3

4

5

∆σ (MPa)

0

1

2

3

4

5

200°C

increment after SR at three temperatures

temperature (°C)

The local solute concentration increment, ∆c, on the dislocations can be expressed as [43-45] **Figure 64.** Strain dependence of the stress increment after SR at two temperatures (AZ63 alloy).

Fig. 64: Strain dependence of the stress Fig. 65: Temperature dependence of the stress increment after SR at two SR **Figure 65.** Temperature dependence of the stress increment after SR at two SR starting strains (AZ91 alloy).

starting strains (AZ91 alloy).

]. (16)

σ = σd + σf, (15)

The flow stress, necessary for the dislocation movement, may be expressed as a sum of two components

$$
\sigma = \sigma\_d + \sigma\_{f'} \tag{15}
$$

where*σ* d is the dislocation component due to strong obstacles (e.g. forest dislocations, precipitates, grain boundaries) and *σ*<sup>f</sup> is the friction stress due to the interaction between the solute atoms and moving dislocations. Mulford and Kocks [41] consider that the dislocation component of the flow stress rather than the friction stress is affected by dynamic strain ageing. Balík and Lukáč [42] take into account the influence of solute atoms on the both flow stress components in the dynamic strain ageing regime.

The local solute concentration increment, Δc, on the dislocations can be expressed as [43-45]

$$
\Delta \mathcal{c} = \mathcal{c} - \mathcal{c}\_0 = \Delta \mathcal{c}\_M \left[ 1 - \left( \exp \left( -t\_a / t\_0 \right)^r \right) \right] \tag{16}
$$

Here *c* is the local solute concentration in the dislocation core, *c*<sup>0</sup> is the nominal solute concen‐ tration in the matrix and Δ*cM* is the maximum concentration increment. The exponent *r* is equal 2/3 or 1/3 for bulk or pipe diffusion, respectively. The relaxation time *t*<sup>0</sup> depends on the binding energy between a dislocation and a solute atom, on solute concentration, and on the diffusion coefficient of solute atoms. *t*<sup>0</sup> is inversely proportional to the diffusion coefficient in the case of bulk diffusion, whereas for pipe diffusion 1/*t*0~ *Dρ*<sup>f</sup> 3/2 [46], where *ρ*<sup>f</sup> is the density of forest dislocations.

The mean ageing time *t*<sup>a</sup> is identified with the mean waiting time *t*w of mobile dislocations on the obstacles. *t*w, is connected with the strain rate by the Orowan equation

$$\dot{\varepsilon} = b \, \rho\_m \, \Lambda \, / \, t\_w = b \Omega \, / \, t\_w \tag{17}$$

where*ρ* m is the mobile dislocation density, Λ is the mean free path of dislocations. If it is assumed that the forest dislocations are the rate controlling obstacles then Λ=1/*ρ*<sup>f</sup> 1/2. The elementary plastic strain per activation event *Ω*=*bρ*m*ρ*<sup>f</sup> -1/2 is strain dependent; it may exhibit a local maximum at a certain strain [43, 44].

If we consider that both the dislocation stress component and the friction stress are influenced by solutes, then the flow stress may be decomposed into a non-aged part *σ*na and a dynamic strain ageing part *σ*a. The strain rate and temperature dependences of the stress increment due to dynamic strain ageing are simply related to the kinetic law Δ*c*(*t*w) for the local concentration changes. Hence,

#### Influence of Solute Atoms on Deformation Behaviour of Selected Magnesium Alloys http://dx.doi.org/10.5772/58949 43

$$
\sigma\_a = \left(f\_1 + f\_2\right) \Delta c \left(t\_w\right). \tag{3}
$$

$$
\text{or}
$$

$$
\sigma\_a = \left(f\_1 + f\_2\right) \left[1 - \exp\left(-X\right)\right] \text{(b)}
$$

The first term *f*<sup>1</sup> corresponds to the dislocation-dislocation interaction influenced by dynamic strain ageing, while the second term *f*2 results from the solute atoms-dislocation interaction influenced by dynamic strain ageing. *X*=(*bΩ/ ε*˙ *t*0). Combining the relations (16) – (18), the negative strain rate dependence of the dynamic strain-ageing component follows. This causes the negative strain rate dependence of the yield stress, which is observed at certain tempera‐ tures and in a certain strain rate range. The negative slope in the strain rate dependence of the yield stress was also observed in Mg+0.7Nd alloy deformed at room temperature (Fig. 57). The observed decrease in the flow stress (negative values of Δ*σ*) due to changes from a lower strain rate to a higher one is also the result of dynamic strain ageing.

Solute atoms locking dislocations cause the observed stress increase after stress relaxation, which depends on strain and on temperature. An increase in the flow stress is needed to move the dislocations after the stress relaxation. It is reasonable to assume that Δ*σ* is proportional to the number of impurities on the dislocation line. The density of dislocations increases with strain, while the concentration of solute atoms is constant. Thus, the stress increase, Δ*σ*, after relaxation due to dynamic strain ageing should decreases with strain, which is observed. The value of *Ω* decreases with strain [44, 45] and hence, *t*w also decreases. This leads to the observed decrease in Δ*σ*. It should be mentioned that the post relaxation behaviour and the values of Δ*σ* depend also on the time of relaxation [47].


#### **8. Nomenclature**

The flow stress, necessary for the dislocation movement, may be expressed as a sum of two

, *d f*

 s

where*σ* d is the dislocation component due to strong obstacles (e.g. forest dislocations, precipitates, grain boundaries) and *σ*<sup>f</sup> is the friction stress due to the interaction between the solute atoms and moving dislocations. Mulford and Kocks [41] consider that the dislocation component of the flow stress rather than the friction stress is affected by dynamic strain ageing. Balík and Lukáč [42] take into account the influence of solute atoms on the both flow stress

The local solute concentration increment, Δc, on the dislocations can be expressed as [43-45]

Here *c* is the local solute concentration in the dislocation core, *c*<sup>0</sup> is the nominal solute concen‐ tration in the matrix and Δ*cM* is the maximum concentration increment. The exponent *r* is equal 2/3 or 1/3 for bulk or pipe diffusion, respectively. The relaxation time *t*<sup>0</sup> depends on the binding energy between a dislocation and a solute atom, on solute concentration, and on the diffusion coefficient of solute atoms. *t*<sup>0</sup> is inversely proportional to the diffusion coefficient in the case

The mean ageing time *t*<sup>a</sup> is identified with the mean waiting time *t*w of mobile dislocations on

/ / *mw w*

assumed that the forest dislocations are the rate controlling obstacles then Λ=1/*ρ*<sup>f</sup>

where*ρ* m is the mobile dislocation density, Λ is the mean free path of dislocations. If it is

If we consider that both the dislocation stress component and the friction stress are influenced by solutes, then the flow stress may be decomposed into a non-aged part *σ*na and a dynamic strain ageing part *σ*a. The strain rate and temperature dependences of the stress increment due to dynamic strain ageing are simply related to the kinetic law Δ*c*(*t*w) for the local concentration

the obstacles. *t*w, is connected with the strain rate by the Orowan equation

e r

( ) 0 0 1 exp / . *<sup>r</sup> M a* D = - =D - - *cc c c t t* é ù

= + (15)

ê ú ë û (16)

& *b tbt* L= = W (17)

3/2 [46], where *ρ*<sup>f</sup> is the density of forest


1/2. The

ss

components in the dynamic strain ageing regime.

42 Magnesium Alloys - Properties in Solid and Liquid States

of bulk diffusion, whereas for pipe diffusion 1/*t*0~ *Dρ*<sup>f</sup>

elementary plastic strain per activation event *Ω*=*bρ*m*ρ*<sup>f</sup>

local maximum at a certain strain [43, 44].

components

dislocations.

changes. Hence,


#### **Acknowledgements**

Z.T. and P.L. are grateful for the financial support of the Czech Science Foundation (project P204/12/1360). P.P. is grateful for the financial support to the Slovak Grant Agency for Science (project VEGA No. 1/0797/12).

#### **Author details**

Zuzanka Trojanová1 , Peter Palček<sup>2</sup> , Pavel Lukáč<sup>1</sup> and Zdeněk Drozd<sup>1</sup>

1 Charles University in Prague, Faculty of Mathematics and Physics, Czech Republic

2 University of Žilina, Faculty of Mechanical Engineering, Slovak Republic

#### **References**


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**Quantities (selected) Alloys**

Z.T. and P.L. are grateful for the financial support of the Czech Science Foundation (project P204/12/1360). P.P. is grateful for the financial support to the Slovak Grant Agency for Science

and Zdeněk Drozd<sup>1</sup>

, Pavel Lukáč<sup>1</sup>

1 Charles University in Prague, Faculty of Mathematics and Physics, Czech Republic

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 total dislocation density AJ91 Mg-9wt%Al-1wt%Sr k Boltzmann constant AJ62 Mg-6wt%Al-2wt%Sr c atomic concentration of solute atoms AX41 Mg-4wt%Al-1wt%Ca CRSS critical resolved shear stress (τ0) AX62 Mg-6wt%Al-2wt%Ca CYS compression yield stress (σ02) AX91 Mg-9wt%Al-1wt%Ca MCS maximum compression strength (σmax) AZ63 Mg-6wt%Al-3wt%Zn TYS tensile yield stress (σ02) AZ91 Mg-9wt%Al-1wt.%Zn MTS maximum tensile strength (σmax) AE42 Mg-4wt%Al-2wt%RE SR stress relaxation ZE41 Mg-4wt%Zn-1wt%RE SEM scanning electron micrograph QE22 Mg-2wt%Ag-2wt%RE

44 Magnesium Alloys - Properties in Solid and Liquid States

ρt

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**Texture, Microstructure, and Mechanical Properties of Calcium-Containing Flame-Resistant Magnesium Alloy Sheets Produced by Twin-Roll Casting and Sequential Warm Rolling**

Masafumi Noda, Tomomi Ito, Yoshio Gonda, Hisashi Mori and Kunio Funami

Additional information is available at the end of the chapter

http://dx.doi.org/10.5772/58940

#### **1. Introduction**

Magnesium (Mg) alloys are attracting attention as metallic materials of the next generation because of their good specific strength, specific rigidity, earthquake resistance, and machina‐ bility, as well as the abundance of resources available for their production [1]. Because of their lightness, Mg alloys are expected to be suitable materials for replacing aluminum alloys in automobiles and in rail and aerospace transportation devices [2,3]. Research has been per‐ formed on improving the mechanical properties, corrosion resistance, and workability of Mg alloys [4,5]. Hot working is required in plastic forming of Mg alloys because of their crystal structure [6]. Problems associated with Mg alloys include their high production costs, their high flammability, and the marked effects of additive elements on the various properties of the alloys. Most structural parts currently manufactured from Mg alloys are die cast, Thixo‐ molded [7,8], semicontinuously cast, or gravity cast [8,9], or they are forged from such materials. The amount of wrought materials that are used is low, representing less than 10% of all Mg alloys that are used [10]. Recently, the problems associated with Mg alloys have been resolved by increasing their strength [11–13] and by improving their creep properties [14], heat resistance [15], and formability [16,17] by adding trace amounts of various elements [18,19], by age-hardening effects [20], by crystal-grain refinement [21,22], and by texture control [23,24]. However, there is still a need for incombustible or noncombustible Mg alloys [25-27] for use in structural or construction components. Flame-resistant and noncombustible Mg alloys have recently been developed by adding Ca to Mg alloys [25–28], and these materials

© 2014 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

have received flame-resistance and noncombustion certifications in the rail, aerospace, and construction sectors.

There are many choices of starting material for manufacturing sheet products, and use of twinroll cast (TRC) materials [29–32] should be examined if production costs are a consideration. However, TRC materials have a number of problems, such as solidifying segregation at the slab thickness center, intrusion of inclusions during casting, and restrictions on casting conditions, such as the casting length, liquid pressure, rolling speed, and quenching capacity [30–32]. Moreover, most generic TRC materials in use are Al alloys or AZ-type Mg alloys that do not preferentially form intermetallic compounds [30–33]. TRC materials described in the literature are subjected to a short TRC process, so that the melt is exposed to the air for only a short time and there is no blackening of the sheet surface through ignition during casting. There is a report that the manufacture of a sheet material from AZ-type Mg alloys can be problematic unless the twin-rolling speed is unusually slow at casting temperatures exceeding 620 °C [34], but there has been no discussion on the effects of the metal texture and the role of plastic deformation after casting. In addition, there has been insufficient discussion on the effect of melt purification before TRC casting and on the role of additive elements on twin-roll casting conditions and the internal texture. Furthermore, there is no report of any comparison with semicontinuous casting in the manufacture of sheet material. We have therefore describe our investigations on the effects of melt purification during melting before TRC on the production of sheet materials, as well as the microstructure and mechanical properties of twinroll cast Ca-containing flame-resistant Mg alloys; we also provide an analysis of the effects on the cast sheets and their durability caused by precipitation of intermetallic compounds formed through addition of Ca.

#### **2. Experimental procedures**

To investigate the effects of the alloy composition on materials produced under TRC condi‐ tions, samples of AZ61 alloy and of AZX611 and AZX612 alloys containing 1 or 2 mass% of added Ca, respectively, were prepared from AM60B Mg and Mg–30%Ca master alloy, pure Al ingots (99.7%), and pure Zn ingots (99.9%). The metals were weighted to the required stoichiometry, placed in a steel crucible, and melted by heating to 700 °C in an electric furnace. Melting was carried out under an argon atmosphere, and argon gas was bubbled through the melt for 20 min. The slug was removed and immediately tapped when the melt temperature reached 660 °C. To prevent combustion of the weighed metals, the melt was isolated from the atmosphere by placing a steel lid on the crucible. The conditions for twin-roll casting were as follows. The speed of both the upper and lower rolls was 20–25 m min–1, and the roll gap was 2 – 3 mm. The rolls used were metal rolls 300 mm wide with a diameter of 200 mm. The surface of the roll was heated to a temperature of more than 100 °C before twin-roll casting prior to decreasing the cooling rate. Grooves were cut into the rolls to facilitate peeling of sheets during twin-roll casting. Figure 1 is a schematic showing the twin-roll casting process. After twin-roll casting, the sheets were left to cool to room temperature and then subjected to strip processing. Sheets measuring 150 mm in width by 200 mm in length were cut from the TRC material for strip processing, with the width direction of the cut sheets aligned in the casting direction. Strip processing was performed by using a two-high rolling mill with 300 mm diameter rolls. The thickness of the original sheet was reduced by 65% in one pass: the roll surface temperature was 245 °C, the workpiece was heated for 10 min at 200–400 °C, the roll speed during processing was 5–25 m min–1, and the workpiece was cooled with water after the strip processing. Tensiletest specimens with a gauge length of 30 mm and a width of 4 mm were prepared from the strip-processed materials, and tensile tests were carried out in air at an initial strain rate of 1.3 × 10–3 s–1. The microstructures of the TRC materials and the strip-processed materials were examined by optical microscopy (OM), and scanning electron microscopy (SEM), and the crystal orientation was analyzed by using electron backscattered diffraction (EBSD).

**Figure 1.** Schematic of a twin-roll caster for strips.

have received flame-resistance and noncombustion certifications in the rail, aerospace, and

There are many choices of starting material for manufacturing sheet products, and use of twinroll cast (TRC) materials [29–32] should be examined if production costs are a consideration. However, TRC materials have a number of problems, such as solidifying segregation at the slab thickness center, intrusion of inclusions during casting, and restrictions on casting conditions, such as the casting length, liquid pressure, rolling speed, and quenching capacity [30–32]. Moreover, most generic TRC materials in use are Al alloys or AZ-type Mg alloys that do not preferentially form intermetallic compounds [30–33]. TRC materials described in the literature are subjected to a short TRC process, so that the melt is exposed to the air for only a short time and there is no blackening of the sheet surface through ignition during casting. There is a report that the manufacture of a sheet material from AZ-type Mg alloys can be problematic unless the twin-rolling speed is unusually slow at casting temperatures exceeding 620 °C [34], but there has been no discussion on the effects of the metal texture and the role of plastic deformation after casting. In addition, there has been insufficient discussion on the effect of melt purification before TRC casting and on the role of additive elements on twin-roll casting conditions and the internal texture. Furthermore, there is no report of any comparison with semicontinuous casting in the manufacture of sheet material. We have therefore describe our investigations on the effects of melt purification during melting before TRC on the production of sheet materials, as well as the microstructure and mechanical properties of twinroll cast Ca-containing flame-resistant Mg alloys; we also provide an analysis of the effects on the cast sheets and their durability caused by precipitation of intermetallic compounds formed

To investigate the effects of the alloy composition on materials produced under TRC condi‐ tions, samples of AZ61 alloy and of AZX611 and AZX612 alloys containing 1 or 2 mass% of added Ca, respectively, were prepared from AM60B Mg and Mg–30%Ca master alloy, pure Al ingots (99.7%), and pure Zn ingots (99.9%). The metals were weighted to the required stoichiometry, placed in a steel crucible, and melted by heating to 700 °C in an electric furnace. Melting was carried out under an argon atmosphere, and argon gas was bubbled through the melt for 20 min. The slug was removed and immediately tapped when the melt temperature reached 660 °C. To prevent combustion of the weighed metals, the melt was isolated from the atmosphere by placing a steel lid on the crucible. The conditions for twin-roll casting were as follows. The speed of both the upper and lower rolls was 20–25 m min–1, and the roll gap was 2 – 3 mm. The rolls used were metal rolls 300 mm wide with a diameter of 200 mm. The surface of the roll was heated to a temperature of more than 100 °C before twin-roll casting prior to decreasing the cooling rate. Grooves were cut into the rolls to facilitate peeling of sheets during twin-roll casting. Figure 1 is a schematic showing the twin-roll casting process. After twin-roll casting, the sheets were left to cool to room temperature and then subjected to strip processing. Sheets measuring 150 mm in width by 200 mm in length were cut from the TRC material for

construction sectors.

50 Magnesium Alloys - Properties in Solid and Liquid States

through addition of Ca.

**2. Experimental procedures**

**Figure 2.** Surface of cast material plate with ignition (a) and without ignition (c), and optical micrographs of twin-roll cast material with ignition (b) and without ignition (d).

**Figure 3.** Nominal stress–strain curves for AZ61 and AZX611 twin-roll cast materials.

**Figure 4.** Cross-sectional optical micrographs [(a), (c), and (e)] and SEM micrographs of fracture surfaces [(b), (d), and (f)] of AZ61 [(a)–(d)] and AZX611 [(e) and (f)] twin-roll cast materials with ignition during heating [(a) and (b)] and without ignition during heating [(c)–(f)].

#### **3. Results and discussions**

**Figure 3.** Nominal stress–strain curves for AZ61 and AZX611 twin-roll cast materials.

52 Magnesium Alloys - Properties in Solid and Liquid States

**Figure 4.** Cross-sectional optical micrographs [(a), (c), and (e)] and SEM micrographs of fracture surfaces [(b), (d), and (f)] of AZ61 [(a)–(d)] and AZX611 [(e) and (f)] twin-roll cast materials with ignition during heating [(a) and (b)] and

without ignition during heating [(c)–(f)].

#### **3.1. The twin-roll casting process and microstructure**

Figure 2(a)–(d) show external views and internal microstructures of sheets subjected to melting in an argon (Ar) atmosphere, bubbling with argon gas, and twin-roll casting. Slight combustion caused oxides to form in the melt and increased the frequency of intrusion of oxides and impurities. Figure 2(a) and 2(b) show the presence of black spots on the surface of the sheet and the inclusion of oxides and impurities in the microstructure. The nominal stress–strain curves of the TRC AZ61 and AZX611 alloys are shown in Figure 3, and the vertical crosssections and fracture surface microstructure are shown in Figures 4(a)–(f). Two types of AZ61 alloys are shown in Figure 4(a)–(d), one in which combustion occurred during melting and one in which it did not occur. The AZX611 alloy melt did not undergo combustion during TRC. When no combustion occurred during melting, the yield stress (YS) was 116 MPa, the ultimate tensile strength (UTS) was 239 MPa, and the elongation was 19% [Figures 4(a) and 4(b)], whereas when combustion occurred, the YS was 82 MPa, the UTS was 180 MPa, and the elongation was 13% [Figures 4(c) and 4(d)]. Intrusion of inclusions and oxides caused early fracture without sufficient plastic deformation [35], and the tensile properties were stable when melts that had not undergone combustion were used. Figure 4 shows that intrusion of oxides caused by combustion results in propagation of cracks, formation of voids, and the appearance of partial brittle fracture. Most of the brittle fracture occurs near the surface of the TRC material, implying that the solidification rate is faster at the surface of the sheet than in its interior, that internal latent heat causes coarsening of the microstructure, and that crack propagation along grain boundaries is initiated from internal voids, resulting ultimately in fracture. In contrast, AZ61 alloy that had not undergone combustion showed a standard ductile fracture surface. In AZX611 alloy containing 1 mass% of Ca, no intrusions of oxides were found and the YS (139 MPa) was higher than that of AZ61 alloy; however, the elongation was only about 8%. Because Al–Ca compounds form at grain boundaries, as shown in Figures 4(e) and 4(f), the reason for early fracture compared with AZ61 alloy is likely to be inhibition of plastic deformation by Al–Ca compounds.

Figure 5 shows the SM and OM microstructures of TRC AZ61 and AZX611 alloys for the case where the roll was at room temperature, the melt temperature was 660 °C, and the roll speed was 25 m min–1. Figures 5(a) and 5(b) show that no cracks were present at the surface of the TRC material in the AZ61 alloy, but that cracks appeared at grooves in the AZX611 alloy. Al– Ca compounds are known to form in AZX611 alloy [3,11]. Figures 5(c) and 5(d) show the OM microstructure after removal of 0.1 mm of material from the sheet surface by mechanical polishing, and they show that Al–Ca compounds formed near cracks in the AZX611 alloy. There was considerable segregation near the surface in the processed sheets shown in Figure 5(d), suggesting that the cooling rate in twin-roll casting is too rapid for alloy systems that form compounds.

Figure 6 shows an example in which Al–Ca compounds segregate at the surface of the TRC material. Because Al–Ca compounds form where the quenching capacity is largest, the rate of solidification of the melt in the twin-roll casting process was reduced by adjusting the

**Figure 5.** Optical micrographs of the surfaces of twin-roll cast strips of AZ61 [(a), (b)] and AZX611 [(c), (d)].

**Figure 6.** Example showing the segregation of Al–Ca compounds in a rapidly cooled region of twin-roll cast material.

temperature of the roll surface to 100 °C and decreasing the roll speed to 20 m min–1. Optical micrographs showing the microstructure of the surface of the resulting sheet and of a cross section perpendicular to the surface are presented in Figure 7. No cracks were found and no segregation of Al–Ca compounds was observed in the prepared sheet. In this case, the sheet temperature immediately after twin-roll casting was 560 °C. By reducing the cooling rate from that of fast solidification at over 100 °C s–1, characteristic of twin-roll casting [29,33], to around 50 °C s–1, samples of TRC Mg alloy material 300 mm wide by 5 m long that form intermetallic compounds can be prepared, as shown in Figure 8.

**Figure 7.** Optical micrographs of AZX611 twin-roll cast material solidified at a lower rate (rise to roll-surface tempera‐ ture, lower roll-mill speed). Observations were made from the direction of the surface (a) and in the perpendicular direction (b).

#### **3.2. Comparisons of the texture, microstructure, and mechanical properties of twin-roll cast material with those of normally cast material**

temperature of the roll surface to 100 °C and decreasing the roll speed to 20 m min–1. Optical micrographs showing the microstructure of the surface of the resulting sheet and of a cross section perpendicular to the surface are presented in Figure 7. No cracks were found and no segregation of Al–Ca compounds was observed in the prepared sheet. In this case, the sheet temperature immediately after twin-roll casting was 560 °C. By reducing the cooling rate from that of fast solidification at over 100 °C s–1, characteristic of twin-roll casting [29,33], to around 50 °C s–1, samples of TRC Mg alloy material 300 mm wide by 5 m long that form intermetallic

**Figure 6.** Example showing the segregation of Al–Ca compounds in a rapidly cooled region of twin-roll cast material.

**Figure 5.** Optical micrographs of the surfaces of twin-roll cast strips of AZ61 [(a), (b)] and AZX611 [(c), (d)].

compounds can be prepared, as shown in Figure 8.

54 Magnesium Alloys - Properties in Solid and Liquid States

The thermal conductivities of AZ61, AZX611, and AZX612 alloy are 63.6, 70.2, and 77.2 W / mK, respectively; therefore, the thermal conductivity increases with addition of Ca. This means that the optimal twin-roll casting conditions will be different for each alloy and, for each sheet, the surface and the interior will have different microstructures. Figure 9 shows the surface and interior OM microstructures of AZ61 and AZX611 TRC alloys. The grain size at the surface of the TRC material is about 2.5 times larger than that in the interior. The surface microstructure consists of coarse grains 70–100 μm in size or dendrites, whereas the interior consists of refined equiaxial grains 25–40 μm in size. The region with refined equiaxial grains account for twothirds of the thickness of the sheet.

The inverse pole figure (IPF) map and the pole figure (PF) map obtained by EBSD analysis of the surface and interior of TRC AZX611 alloy are shown in Figure 10. TRC material does not show any specific crystal orientation in either the surface or the interior, regardless of grain size. As shown in Figure 9(a), the texture orientation is high (10.9) because of the large size of the crystal grains compared with the visual area for measurement. The microstructure orientation of the interior texture is 3.8, indicating a random orientation. One noteworthy point is that the OM microstructure and IPF map suggest that a grain boundary forms between the refined-grain region and coarse-grain region. However, measurement of the crystal orientation between adjacent grains showed that coarse grains adjacent to regions of refined grains are in fact agglomerates of refined subgrains. This means that differences between the surface and interior microstructure can be reduced by optimizing the twin-roll casting method or by a cooling and homogenization processes after twin-roll casting. No cracks at interfaces or grain boundaries were found in either the coarsegrain region or the refined-grain region.

Table 1 lists the average grain sizes at the surfaces and in the interiors of AZ61, AZX611, and AZX612 TRC alloys, as well as the proportion of compound formation and the hardness. These measurements were performed by using 3 mm thick TRC sheets. As the Ca content increases, the crystal grain size becomes smaller, the area ratio of intermetallic compounds increases, and the hardness increases. Figure 11 shows the OM structure and the IPF and FP maps for AZX611 alloy semicontinuously cast at a cooling rate of 25 °C s–1. The average grain size of the standard gravity-cast material is around 600–800 μm [11]. This can be refined to about 100 μm by increasing the cooling rate, thereby making possible refined dispersion of Al–Ca compounds in the Mg phase [35]. Here, the texture orientation is 10.9 and a dendrite microstructure is present in the crystal texture, showing the same trends as the microstructure of the surface of TRC material. The YS and UTS of semicontinuously cast material have been reported to be 90 MPa and 155 MPa, respectively, and the elongation is 5%. The elongation is lower than that found in tensile tests on TRC AZX611 material, although the tensile behavior is similar. Therefore, large changes in single uniaxial tensile tests were not observed when the cooling rate exceeded a certain value and the average grain size after casting was less than 100 μm, Addition of Ca causes minute cracks to form more easily, as mentioned before, and similar results are found in semicontinuously cast material. On the other hand, as Al–Ca compounds are smaller and undergo refined dispersion in the Mg phase in TRC material, as shown in Figures 10 and 11, the rolling reduction per pass can be large in the strip-processing technique described later. The results agree with the maximum rolling reduction curve for extruded AZX311 determined by Noda *et al.* [11] The use of extruded materials or TRC materials instead of cast materials with a dendrite microstructure and/or coarse grains results in good stripprocessing performance as a result of random crystal orientation, formation of a refined texture, and refined dispersion of Al–Ca compounds.

Texture, Microstructure, and Mechanical Properties of Calcium-Containing ... http://dx.doi.org/10.5772/58940 57

equiaxial grains 25–40 μm in size. The region with refined equiaxial grains account for two-

The inverse pole figure (IPF) map and the pole figure (PF) map obtained by EBSD analysis of the surface and interior of TRC AZX611 alloy are shown in Figure 10. TRC material does not show any specific crystal orientation in either the surface or the interior, regardless of grain size. As shown in Figure 9(a), the texture orientation is high (10.9) because of the large size of the crystal grains compared with the visual area for measurement. The microstructure orientation of the interior texture is 3.8, indicating a random orientation. One noteworthy point is that the OM microstructure and IPF map suggest that a grain boundary forms between the refined-grain region and coarse-grain region. However, measurement of the crystal orientation between adjacent grains showed that coarse grains adjacent to regions of refined grains are in fact agglomerates of refined subgrains. This means that differences between the surface and interior microstructure can be reduced by optimizing the twin-roll casting method or by a cooling and homogenization processes after twin-roll casting. No cracks at interfaces or grain boundaries were found in either the coarse-

Table 1 lists the average grain sizes at the surfaces and in the interiors of AZ61, AZX611, and AZX612 TRC alloys, as well as the proportion of compound formation and the hardness. These measurements were performed by using 3 mm thick TRC sheets. As the Ca content increases, the crystal grain size becomes smaller, the area ratio of intermetallic compounds increases, and the hardness increases. Figure 11 shows the OM structure and the IPF and FP maps for AZX611 alloy semicontinuously cast at a cooling rate of 25 °C s–1. The average grain size of the standard gravity-cast material is around 600–800 μm [11]. This can be refined to about 100 μm by increasing the cooling rate, thereby making possible refined dispersion of Al–Ca compounds in the Mg phase [35]. Here, the texture orientation is 10.9 and a dendrite microstructure is present in the crystal texture, showing the same trends as the microstructure of the surface of TRC material. The YS and UTS of semicontinuously cast material have been reported to be 90 MPa and 155 MPa, respectively, and the elongation is 5%. The elongation is lower than that found in tensile tests on TRC AZX611 material, although the tensile behavior is similar. Therefore, large changes in single uniaxial tensile tests were not observed when the cooling rate exceeded a certain value and the average grain size after casting was less than 100 μm, Addition of Ca causes minute cracks to form more easily, as mentioned before, and similar results are found in semicontinuously cast material. On the other hand, as Al–Ca compounds are smaller and undergo refined dispersion in the Mg phase in TRC material, as shown in Figures 10 and 11, the rolling reduction per pass can be large in the strip-processing technique described later. The results agree with the maximum rolling reduction curve for extruded AZX311 determined by Noda *et al.* [11] The use of extruded materials or TRC materials instead of cast materials with a dendrite microstructure and/or coarse grains results in good stripprocessing performance as a result of random crystal orientation, formation of a refined

thirds of the thickness of the sheet.

56 Magnesium Alloys - Properties in Solid and Liquid States

grain region or the refined-grain region.

texture, and refined dispersion of Al–Ca compounds.

**Figure 9.** Optical micrographs of twin-roll cast strips of AZ61 [(a) and (b)] and AZX611 [(c) and (d)] showing the surfa‐ ces of the strips [(a) and (c)] and the interiors of the strips [(b) and (d)].

**Figure 10.** Inverse pole figure (IPF) maps and pole figure (PF) maps of AZX611 twin-roll cast material. The intensity of texture is indicated in the PF maps. Figures (b) and (c) were cropped from the IPF map; (b) shows the surface region and (c) shows the interior.


**Table 1.** Grain size, Vickers hardness, and area frequency of compounds for various twin-roll cast materials.

**Figure 11.** Optical micrograph and IPF and PF maps of AZX611 antigravity suction-cast material cooled at 25 °C s–1. The intensity of texture is indicated in the PF map.

**Figure 12.** Relationship between tensile properties and roll-mill speed for AZ61, AZX611, and AZX612 twin-roll cast materials subjected to a single-pass rolling process.

**Figure 13.** (a) Relationship between the grain size and the annealing temperature for single-pass rolled samples of AZ61, AZX611, AZX 612; and (b) optical micrographs of AZ61, AZX611, and AZX612 materials subjected to single-pass rolling at a sample temperature of 200 °C. The roll-mill speeds are indicated in the optical micrographs.

#### **3.3. Mechanical properties and microstructures of rolled materials**

**Table 1.** Grain size, Vickers hardness, and area frequency of compounds for various twin-roll cast materials.

**Figure 11.** Optical micrograph and IPF and PF maps of AZX611 antigravity suction-cast material cooled at 25 °C s–1.

**Figure 12.** Relationship between tensile properties and roll-mill speed for AZ61, AZX611, and AZX612 twin-roll cast

The intensity of texture is indicated in the PF map.

58 Magnesium Alloys - Properties in Solid and Liquid States

materials subjected to a single-pass rolling process.

Figure 12 shows the tensile properties of TRC AZ61, AZX611, and AZX612 alloys subject‐ ed to strip processing with large one-pass rolling reductions with a strip roll-surface temperature of 245 °C, a specimen temperature of 200 °C, and various roll speeds. The OM microstructures and recrystallized grain sizes are given in Figure 13. The rolling reduc‐ tion was 65% of the original thickness. According to Figure 12, the YS of AZ61alloy decreased monotonously from 260 MPa to 230 MPa, and that of AZX611 decreased monotonously from 310 MPa to 270 MPa when the strip processing speed was increased to 15 m min–1. Increasing the roll speed further to 25 m min–1 resulted in a decrease in the YS to 190 MPa for AZ61 and to 250 MPa for AZX611. The grain size of strip-processed AZ61 and AZX611 alloys was 3 μm regardless of the strip-processing speed. The grain size normally coarsens with increasing strip-processing speed; however, unlike the case of multipass strip processing, the rolling speed has no apparent effect on the crystal texture in the one-pass TRC process because of the effects of processing heating. The Al–Ca

compounds that form at the grain boundaries in TRC AZX611 alloy reorient parallel to the rolling direction after strip processing. On the other hand, the microstructure of process‐ ed AZX612 alloy is similar to that of AZ61 and AZX611 alloys, as described above, but the tensile properties are independent of the strip-processing speed. The area ratio of Al–Ca compounds in TRC AZX611 alloy is 7.1%. However, in AZX612 alloy, this value increas‐ es to 18.1%, and the decrease in strength for strips processed with a large reduction ratio is small. Furthermore, as shown in Figure 13, Al–Ca compounds dispersed in the Mg phase inhibit the growth of Mg phase grains. Solution treatment or homogenizing treatment is typically necessary when Al–Ca compounds develop. It is easier to strip process AZX612 alloy than gravity-cast material, and the dispersion of compounds is the dominant factor in plastic deformation. Figure 14 shows the relationship between the annealing tempera‐ ture and the grain size when a material strip processed at a roll speed of 5 m min–1 is annealed for one hour; it also shows the OM microstructure of material annealed for one hour at 400 °C. Crystal grains of AZ61 alloy grew rapidly at an annealing temperature of 350 °C. However, as the amount of added Ca was increased, the ultimate grain size was 10 μm in AZX611 alloy and 7 μm in AZX612 alloy, even when the annealing tempera‐ ture was increased. There are no compounds that inhibit grain growth in AZ61, whereas Al–Ca compounds inhibit grain growth in AZX611 and AZX612 alloys. The area ratio of Al–Ca compounds is higher in AZX612, and large amounts of dispersion into the Mg phase after strip processing effectively inhibit grain growth. These results suggest that adding more Ca results in an improvement in thermal resistivity, and agrees with a report on improvement of creep-resistance properties [14].

**Figure 14.** (a) Relationship between the annealing temperature and the grain size for AZ61, AZX611, and AZX612 single-pass-rolled materials. Annealing was performed at 200, 300, 350, or 400 °C for one hour. (b) Optical micro‐ graphs of materials annealed at 400 °C for one hour.

The TRC material prepared in this research had an equiaxial microstructure, permitting a large rolling reduction of about 65% per pass. For Mg alloys that form Al–Ca compounds on addition of Ca, resulting in deterioration of the plastic-deformation performance, twinroll casting is an effective processing method, although it is limited to the manufacture of sheet materials. With regard to the microstructure after strip processing, although the internal microstructure of the TRC material varied with the twin-roll casting conditions, in no case was there mixing of grain size in the internal microstructure, that is, there was no mixing of the coarse microstructure at the surface of the sheets with the refined internal microstructure. No solution treatment or microstructure homogenization treatment is carried out after twin-roll casting of AZ61 alloy; as a result, the β-phase remains elongat‐ ed along the rolling direction.

#### **3.4. Corrosion behavior of rolled magnesium alloys**

compounds that form at the grain boundaries in TRC AZX611 alloy reorient parallel to the rolling direction after strip processing. On the other hand, the microstructure of process‐ ed AZX612 alloy is similar to that of AZ61 and AZX611 alloys, as described above, but the tensile properties are independent of the strip-processing speed. The area ratio of Al–Ca compounds in TRC AZX611 alloy is 7.1%. However, in AZX612 alloy, this value increas‐ es to 18.1%, and the decrease in strength for strips processed with a large reduction ratio is small. Furthermore, as shown in Figure 13, Al–Ca compounds dispersed in the Mg phase inhibit the growth of Mg phase grains. Solution treatment or homogenizing treatment is typically necessary when Al–Ca compounds develop. It is easier to strip process AZX612 alloy than gravity-cast material, and the dispersion of compounds is the dominant factor in plastic deformation. Figure 14 shows the relationship between the annealing tempera‐ ture and the grain size when a material strip processed at a roll speed of 5 m min–1 is annealed for one hour; it also shows the OM microstructure of material annealed for one hour at 400 °C. Crystal grains of AZ61 alloy grew rapidly at an annealing temperature of 350 °C. However, as the amount of added Ca was increased, the ultimate grain size was 10 μm in AZX611 alloy and 7 μm in AZX612 alloy, even when the annealing tempera‐ ture was increased. There are no compounds that inhibit grain growth in AZ61, whereas Al–Ca compounds inhibit grain growth in AZX611 and AZX612 alloys. The area ratio of Al–Ca compounds is higher in AZX612, and large amounts of dispersion into the Mg phase after strip processing effectively inhibit grain growth. These results suggest that adding more Ca results in an improvement in thermal resistivity, and agrees with a report on

**Figure 14.** (a) Relationship between the annealing temperature and the grain size for AZ61, AZX611, and AZX612 single-pass-rolled materials. Annealing was performed at 200, 300, 350, or 400 °C for one hour. (b) Optical micro‐

improvement of creep-resistance properties [14].

60 Magnesium Alloys - Properties in Solid and Liquid States

graphs of materials annealed at 400 °C for one hour.

There are many reports on the corrosion resistance of generic AZ-type Mg alloys, but few reports on those of noncombustible Ca-containing Mg alloys. We therefore examined the corrosion properties of such alloys.

Samples of TRC Mg–3Al–1Zn–1Ca (AZX311), Mg–6Al–1Zn–1Ca (AZX611), and Mg–10Al– 0.2Mn–1Ca (AMX1001) alloy sheets were warm-forged from a thickness of 3 mm to 1 mm. Although TRC materials have good strip-processing properties, to avoid mixing of grains in the internal microstructure as a result of strip processing, three-pass processing was employed with a roll speed of 10 m min–1, a roll-surface temperature of 200 °C, and a specimen temper‐ ature of 250 °C. Reheating for 1 min was carried out between each pass, and the specimens were water cooled at the end of the process. Immersion tests were carried out in 5% aqueous NaCl at a constant temperature of 25 °C for up to 120 hours. Changes in weight and in surface microstructure were observed. After immersion, the material was prepared for examination by means of surface polishing, degreasing with acetone, and cutting into samples 20 mm long by 20 mm wide.

The average grain size of AZX311, AZX611, and AMX1001 strip-rolled alloys was 3.5 μm, and refined equiaxial grains and some elongated microstructure that had not recrystallized remained. The YS of all the alloys was around 330 MPa, the UTS was 350 MPa, and the elongation was 5% [3, 11, 35]. The relationship between the change in weight and the immer‐ sion time shown in Figure 15(a) indicates that, for noncombustible Mg alloy containing 1 mass % of Ca, an increase in the Al content reduces the change in weight with increasing immersion time in a manner similar to that shown by generic AZ-type Mg alloys. The weight monoto‐ nously decreased for immersion times of up to 100 hours, and then changed quadratically regardless of the Al content. In the region where the weight decreased monotonously, filiform corrosion and partial pitting was observed at the sheet surface, as shown in the OM micro‐ structure of Figure 15(b); however, these processes did not lead to significant weight losses because the crystal microstructure was refined. On the other hand, filiform corrosion and pitting per unit area increased with time, and these propagated and combined, leading to pitting of the entire surface of the sheet and a significant reduction in weight. Refined disper‐ sion of Ca or Al–Ca compounds in the more-refined Mg phase as a result of casting and strip processing does not result in significant reduction in the corrosion resistance caused by intermetallic compounds if the amount of added Ca is only 1 mass%.

**Figure 15.** (a) Relationship between weight loss and immersion time for AZX311, AZX611, and AMX1001 rolled mate‐ rials. (b) Optical micrographs of plate surfaces after immersion tests in an 5% aqueous NaCl solution.

#### **4. Summary**

In this present study, we investigated on the effects of melt purification during melting before twin-roll casting on the production of sheet metals. Using an Ar gas melting atmosphere prevented combustion at ingot alloy surface, and the amount of slug and oxide inclusions in the cast material decreased. By reducing the cooling rate from that of fast solidification at over 100 °C s–1, characteristic of twin-roll casting, to around 50 °C s–1, samples of TRC Mg alloy material 300 mm wide by 5 m long that form intermetallic compounds can be prepared without cracking. The grain size at the surface of the TRC material is about 2.5 times larger than that in the interior. As the Ca content increases, the crystal grain size becomes smaller and the area ratio of intermetallic compounds increases. The use of TRC materials instead of cast materials with a dendrite microstructure and/or coarse grains results in good strip-processing perform‐ ance as a result of random crystal orientation, formation of a refined microstructure, and refined dispersion of Al–Ca compounds. Refined dispersion of Ca or Al–Ca compounds in the more-refined Mg phase as a result of casting and strip processing does not result in significant reduction in the corrosion resistance.

#### **Author details**

sion of Ca or Al–Ca compounds in the more-refined Mg phase as a result of casting and strip processing does not result in significant reduction in the corrosion resistance caused by

**Figure 15.** (a) Relationship between weight loss and immersion time for AZX311, AZX611, and AMX1001 rolled mate‐

In this present study, we investigated on the effects of melt purification during melting before twin-roll casting on the production of sheet metals. Using an Ar gas melting atmosphere prevented combustion at ingot alloy surface, and the amount of slug and oxide inclusions in the cast material decreased. By reducing the cooling rate from that of fast solidification at over 100 °C s–1, characteristic of twin-roll casting, to around 50 °C s–1, samples of TRC Mg alloy material 300 mm wide by 5 m long that form intermetallic compounds can be prepared without cracking. The grain size at the surface of the TRC material is about 2.5 times larger than that in the interior. As the Ca content increases, the crystal grain size becomes smaller and the area ratio of intermetallic compounds increases. The use of TRC materials instead of cast materials with a dendrite microstructure and/or coarse grains results in good strip-processing perform‐ ance as a result of random crystal orientation, formation of a refined microstructure, and refined dispersion of Al–Ca compounds. Refined dispersion of Ca or Al–Ca compounds in the more-refined Mg phase as a result of casting and strip processing does not result in significant

rials. (b) Optical micrographs of plate surfaces after immersion tests in an 5% aqueous NaCl solution.

**4. Summary**

reduction in the corrosion resistance.

intermetallic compounds if the amount of added Ca is only 1 mass%.

62 Magnesium Alloys - Properties in Solid and Liquid States

Masafumi Noda1 , Tomomi Ito2 , Yoshio Gonda2 , Hisashi Mori3 and Kunio Funami1

\*Address all correspondence to: mk-noda@s7.dion.ne.jp

1 Department of Mechanical Science and Engineering, Chiba Institute of Technology, Narashino, Chiba, Japan

2 Magnesium Division, Gonda Metal Industry Co., Ltd., Sagamihara, Kanagawa, Japan

3 Railway Technical Research Institute, Kokubunji, Tokyo, Japan

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