The Effect of Yttrium Addition on the Microstructure and Mechanical ...

Trans Indian Inst Met (2015) 68(3):331–339 DOI 10.1007/s12666-014-0464-x

REVIEW PAPER

The Effect of Yttrium Addition on the Microstructure and Mechanical Properties of Mg Alloys Lavish Kumar Singh • A. Srinivasan • U. T. S. Pillai • M. A. Joseph • B. C. Pai

Received: 16 June 2014 / Accepted: 10 September 2014 / Published online: 29 October 2014 Ó The Indian Institute of Metals - IIM 2014

Abstract Automotive and aerospace industries shall witness advancement to the next generation if magnesium (Mg) alloys become an integral part of their manufacturing unit primarily because of its light weight. The main limitation which hinders the progress in this direction is the inferior creep properties of Mg alloys. In order to transform this expectation into reality, rare earth (RE) elements are extensively used as alloying elements for improving the room temperature (RT) as well as high temperature (HT) properties. Yttrium (Y) is one of the most extensively used RE elements primarily because of its very high solubility in Mg. Many researchers have studied the influence of Y addition in Mg alloys because of its very high solubility in Mg. However, there is a need to consolidate the work that has been carried out so far which will help in interpolating the future prospects. This review consolidates the work that has been carried out so far in Y addition covering various aspects related to microstructural modifications and RT as well as HT mechanical properties. Keywords AZ91

Yttrium Microstructure Al2Y Mg17Al12

L. K. Singh (&) M. A. Joseph Department of Mechanical Engineering, National Institute of Technology Calicut, Calicut 673601, India e-mail: [email protected] A. Srinivasan U. T. S. Pillai B. C. Pai Materials Science and Technology Division, National Institute for Interdisciplinary Science and Technology (NIIST)-CSIR, Trivandrum 695019, India e-mail: [email protected]

1 Introduction Mg alloys are considered as Green Engineering Materials of the 21st century and possess many advantages, such as light weight, high specific strength, good vibration resistance and electromagnetic shielding properties and have wide application prospects in the field of aerospace, automobile and 3C products such as computer outer shells, camera casings etc. [1]. The extensive uses of Mg alloys in industries are limited because of its inferior HT properties. In order to improve HT properties and to overcome the other disadvantages such as limited workability, poor corrosion resistance and a high degree of shrinkage during solidification, a number of methods such as hot extrusion [2], rapid solidification [3], directional solidification [4] and mechanical alloying [5] have been adopted. But these processing methods have led to increase in production cost. One effective method is alloying, that is, addition of suitable element(s) in minor quantities into Mg. Some of the commonly used alloying elements are Al, Ag, Mn, Zn, Si, Zr and RE elements. In particular, addition of RE elements distinctly influences the microstructure and properties of Mg alloys [6–8]. Y is one of the most frequently used alloying RE element owing to its high solubility (11 wt% at 567 °C) in Mg [9]. Young’s modulus, the lattice parameters and vacancy concentration increases with the increase in Y content in solid solution in Mg matrix [10]. The increase in vacancy with the increase in Y content accounts for the improvement in hardness and tensile properties [11]. Among the Mg based alloys, Mg–Al alloy system has been studied extensively for use in automobiles due to the weight reduction and excellent castability [12, 13]. However, the extensive commercial applications of Mg–Al alloys are limited because of the poor mechanical properties at 120 °C [13–15]. Mg–Zn alloys are one of the

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Trans Indian Inst Met (2015) 68(3):331–339

most promising alloy systems which have moderate strength and corrosion resistance. It responds positively to the addition of a small amount of suitable elements which enhances precipitation and age-hardening behavior [16]. In this review, Mg–Al and Mg–Zn alloy systems are considered separately for studying the changes that occur in their microstructure and mechanical properties as a result of Y addition. The alloy designation along with the nominal chemical compositions used by different authors whose work has been reported in this work is listed in Table 1.

2 Microstructure 2.1 Mg–Al Alloys In general, the microstructure of Mg–Al alloys consists of a-Mg as matrix and b-Mg17Al12 as the secondary phase. Figure 1 shows the microstructure of Mg–6Al alloy which contains a-Mg and Mg17Al12 whereas, the as-cast microstructures of AZ91 alloy consist of a-Mg and eutectic bMg17Al12 ? a-Mg phase [16, 17]. In both the microstructures, Mg17Al12 phase is distributed at the grain boundaries. However, it is continuous in AZ91 whereas it is discontinuous in Mg–6Al alloy. Precipitation of Mg17Al12 phase from the supersaturated a-Mg takes place either in continuous mode or in discontinuous mode. Continuous precipitates form uniformly within the supersaturated a-Mg phase and has lath shaped morphology. Discontinuous precipitates are a lamellar arrangement of Mg17Al12 and Al rich a-Mg phase [18] and their presence at grain boundaries is extremely detrimental for creep properties. Figure 2 shows T6 treated AZ91 alloy with continuous and discontinuous precipitates. It has been reported that certain alloying elements such as RE elements which includes Y can suppress the discontinuous precipitation of Mg17Al12 phase at grain boundaries [19]. When Y is added in Mg–Al alloys, it forms intermetallic only with Al and not with Mg. This is so because larger the electronegativity difference, greater is the tendency to form intermetallic compound and the difference being 0.3 in case of Al and Y and 0.1 for Mg and Y. According to the ternary alloy phase diagram of

Mg–Al–Y, Al and Y can form Al2Y and Al3Y, in which the Y content is about 27 and 32 % respectively [20]. Al2Y is the most frequently encountered intermetallic probably because of its low Y requirement and crystallizing temperature as compared to Al3Y. As the Al–Y intermetallic forms, a significant amount of Al is utilized and therefore, the amount of Al available for the formation Mg17Al12 phase is very less [21]. Besides intermetallic formation, Y addition extensively refines the grains. Wang et al. [17] have reported that the addition of 0.9 % of Y in AZ91 alloy reduces the grain size from 95 to 55 lm. Y acts as a site for heterogeneous nucleation for a-Mg phase because of its structural similarity (hexagonal close-packed), similar lattice parameter and atomic radii with Mg [22]. Sheng et al. [23] explains the reason for the grain refining efficiency of Y in a different manner. The crystallizing temperature of A12Y (980 °C) [20] is much higher than the eutectic reaction temperature of L ? a-Mg ? Mg17Al12 (451 °C), therefore, Al2Y forms first and segregates at the front edge of aMg phase [17]. Thus, constitutional under-cooling appears ahead of the solid–liquid interface [24]. This promotes the nucleation in the under-cooling zone. Meanwhile, in the constitutional undercooling zone, the diffusion of solute occurs at a slow rate which in turn limits the crystal growth rate. Lee et al. [25] have also suggested that addition of solute elements causes constitutional undercooling which is a major driving force for nucleation. Intermetallics may also contribute to nucleation provided the alloying element is a powerful nucleant and the intermetallic is formed before primary phase. Grain refining efficiency of an element can be explained by Growth Restriction Factor (GRF) which conveys the effect of solute on grain refinement. Greater is the GRF, stronger is the refining ability [26]. The GRF is defined as [27]: X GRF ¼ m cðk 1Þ where m is the slope of the liquidus line, c is the initial composition and k is the equilibrium partition coefficient for the element. The growth restriction parameter, m(k - 1), for Y is 1.70, whereas, for tin it is 1.47 which is considered as a

Table 1 Alloy designation along with the respective nominal alloy composition (wt%) Alloy designation

Al

Zn

Mn

Cu

Zr

Mg

AZ91 AZ80

8.60 8.00

0.7 0.6

0.25 0.30

– –

– –

Bal. Bal.

AZ61

6.02

1.01

0.25

–

–

Bal.

AZ31

3.00

1.00

0.30

–

–

Bal.

ZA52

1.67

4.79

–

–

–

Bal.

ZCK630

–

5.64

–

2.92

0.42

Bal.

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Fig. 1 Samples of a Mg-6Al and b AZ91 alloy showing the secondary phase [16, 17]

Fig. 2 Microstructures of T6 treated AZ91 alloy showing continuous and discontinuous precipitates [19]

decent grain refiner for Mg–Al alloys [25]. Zr has the highest GRF value (growth restriction parameter = 38.29) among the common alloying elements [25] yet it is not used as a grain refiner in Mg–Al alloys because the presence of Al effectively suppresses the liquid solubility of Zr in Mg. Also, Zr reacts preferentially with iron present in the melt and forms Fe–Zr intermetallic compound. As a result, the Fe content in the melt reduces and acts as a driving force for rapid uptake of iron from the crucible containing the melt [28]. When Y is added along with other elements in Mg–Al alloys, different intermetallic compounds are formed as shown in Table 2. When Y and Nd are added in Mg–6Al alloy, the compounds such as Al2Y and Al2Nd having high melting point precipitates largely hindering the growth of a-Mg crystal, thereby, refining the grains of Mg–6Al alloy [29]. Moreover, the GRF of Nd is more as compared to Y [30], so the contribution of Nd in grain refinement is more as compared to Y. On the other hand, if the same amount of Y is added to Mg–6Al alloy, the degree of grain refinement observed is less. The average grain size of Mg–6Al alloy with 1.8 % Y addition is 206.8 lm [16] whereas, with 1.8 % RE (1.2 % Y ? 0.6 % Nd) it is 45 lm [31]. When Y is added along with an element having low solubility in Mg, then, the atoms of the element with low solubility are squeezed into the solid–liquid interface during

solidification and these atoms restrain the growth of grains which leads to constitutional supercooling to promote nucleation, thereby, refining the grains [32]. The addition of Y does not always result in grain refinement. If the Y content increases beyond a certain limit, the grains get coarser. It has been reported that in AZ91 alloy, when the Y content increases beyond 0.9 %, the grain coarsening starts [17] and for hot rolled AZ61 the critical limit is 1.4 % Y [33]. Y not only affects the grain size, but also changes the grain morphology. When 0.5 % Sr and 0.5 % Y are added to AZ31 alloy, dendrites appear in the microstructure but when the Y content is increased to 1.0 %, equiaxed grains forms along with small rod shaped phases dispersed at grain boundaries [34]. So, Y addition alters the grain size as well as grain morphology in Mg–Al alloys. 2.2 Mg–Zn Alloys Mg–Zn alloy system has drawn attention owing to its high solid solution strengthening. Zn is approximately three times more effective, on atomic percent basis, than Al in increasing the YS for alloys that are in the solution treated or quenched condition [28]. The phases that are generally present in Mg-Zn alloys are a-Mg and eutectic phase similar to Mg–Al alloys. If Y is added to Mg–Zn–Al alloy

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Table 2 Different intermetallic compounds in Mg–Al alloys as a result of various additive elements along with Yttrium Parent alloy

Alloying elements (%)

Intermetallic compounds

Reference

Mg–6Al

1.2-Y, 0.6-Nd

Al2Y, Al2Nd

[31]

AM50

0.3-Y, 0.6-Ce

Al2Y, Al11Ce3

[32]

AZ31

1.5-Y, 0.5-Sr

Al4Sr, Al2Y, Al3Y

[34]

AZ91

1.0-Y, 1.0-Ca

Al2Y, Al2Ca

[59]

Fig. 3 SEM images of a ZA52Y2 alloy, b ZCK630 ? 0.5Y alloy showing the secondary phases [49, 53]

system, it results in secondary phases like Al11Y3 and Al2Y. When Y is added to Mg–Zn alloy, according to Mg– Zn–Y ternary phase diagram, three kinds of equilibrium phases results. They are W-phase (Mg3Zn3Y2, cubic structure), I-phase (Mg3Zn6Y, icosahedral quasi-crystal structure) and Z-phase (Mg12ZnY) [35–37]. I-phase has attracted much attention because of its high hardness, thermal stability and corrosion resistance along with low coefficient of friction and interfacial energy [38, 39]. Figure 3 shows the secondary phases that are formed in ZA52Y2 and ZCK 630 ? 0.5Y. In Mg–Zn–Y and Mg–Zn– Y–Zr alloys, Zn/Y ratio decides the phases that are present. As reported by Xu et al. [40] in Mg–Zn–Y–Zr alloys, when Zn/Y ratio[4.38, I-phase forms. When the ratio is between 1.10 and 4.38, both I-phase and W-phase are present and when the ratio \1.10, only W-phase is present [41]. But Lee et al. [42] have reported that in Mg–Zn–Y alloys, when Zn/Y ratio \2, there is no formation of I-phase. Table 3 shows the nominal alloy compositions and constituent phases in the as-cast samples identified by X-ray diffraction in Mg–Zn–Y alloys with varying Zn/Y ratio. It is because of these kinds of contradictions that several authors give more emphasis on individual content of Zn and Y rather than Zn/Y ratio. In case of Mg–Zn–Y alloys with low Zn content no binary phase has been detected [43]. This is because the ternary eutectic temperature is higher than that of the binary phase, so at grain boundaries Mg–Zn–Y ternary phases precipitate first. Another study on Mg–Zn–Y with high Zn content has reported that I-phase mainly distributes

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at the triple junction while W-phase distributes at the intergranular grain boundary [44]. However, in case of low Zn content, both the phases are found at the triple junction and intergranular grain boundary [43]. Addition of Y also leads to increase in the eutectic temperature and volume of the secondary phase. From the DTA study, the lowest eutectic temperature of Mg– 5.54Zn–0.74Y–0.39Zr and Mg–5.55Zn–1.72Y–0.37Zr alloys is about 450 °C, which is much higher than the eutectic temperature of Mg–Zn binary alloy (340 °C) [45]. Zhou et al. [46] have also reported that the eutectic temperature of Mg–Zn–Zr alloy increases when Nd and Y are added into it. The volume of the secondary phase increases from 1.79 to 4.26 % as the Y increases from 0.36 to 0.82 wt% in Mg–Zn–Y with low Zn content [43]. In as-cast Mg–Zn–Y alloy, when the wt% of Y are 0.36, 0.82 and 1.54, the volume of secondary phase (%) are 1.79, 4.26 and 4.67 respectively and the respective average grain sizes are 200, 100 and 80 lm respectively [43]. This also reveals the grain refining efficiency of Y in Mg–Zn alloys along with the influence of inermetallics on grain refinement. Similarly, for the as-cast structure of Mg–Zn–Y–Zr alloys with Y content of 0.74, 1.35 and 1.72 %, the average grain size is 80, 64 and 45 lm respectively [47]. This shows extensive grain refining capability of Zr which can be attributed to the fact that its a-allotrope has similar crystal structure and lattice dimensions (for Zr, a = 0.323 nm and c = 0.514 nm and for Mg, a = 0.320 nm and c = 0. 520 nm) as that of Mg [28]. Therefore according to the principle of size and structure matching, Zr acts as a


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Table 3 Nominal alloy compositions and constituent phases in the as-cast samples identified by X-ray diffraction in Mg–Zn–Y alloys with varying Zn/Y ratio [42] Zn (wt%)

Y (wt%)

Mg (wt%)

Zn/Y ratio

Phases

4

0.4

Rem.

10

a-Mg ? Mg7Zn3

8

0.8

Rem.

10

a-Mg ? Mg7Zn3

4

0.6

Rem.

6.7

a-Mg ? I-phase

3

0.5

Rem.

6

a-Mg ? I-phase

3

0.6

Rem.

5

a-Mg ? I-phase

3

1.2

Rem.

2.5

a-Mg ? I-phase ? W-phase

4 3

2.0 1.6

Rem. Rem.

2.0 1.9

a-Mg ? I-phase ? W-phase a-Mg ? W-phase

4

2.2

Rem.

1.8

a-Mg ? W-phase

nucleation core. The refining effect of Y can be mainly due to the following factors: (i) Y changes solution degree of Zn, which decreases the solidus curve and shortens the time for nucleation, and thus reduces the grain size [48] (ii) The formation of Mg–Zn–Y phases, especially the I-phase can effectively restrain the grain growth during the DRX [49] (iii) The low diffusibility of Y contributes to suppression of grain growth [47].

3 Room Temperature Properties 3.1 Mg–Al Alloys In general, addition of Y in Mg–Al alloys increases both the UTS and elongation. Table 4 shows the variation of UTS and elongation with respect to Y addition in different Mg–Al alloys. The increase in UTS is due to grain refinement and formation of stable secondary phase, which exerts a dispersion hardening effect. Further it is reported that an increase in elongation with Y addition is not always possible. The experimental results show that when Y is added in an extruded AZ31 alloy, it is observed that UTS and YS increases but the elongation decreases. The possible reason is the formation of many coarse blocky Al2Y compounds which increase the viscosity of the alloy melt during solidification thus decreasing the flowability and reduces boundary combination strength as it influences micro- crack formation and flaw propagation [23]. It is further said that the addition of Y exerts a positive effect on RT properties only up to a certain limit, beyond which its addition is detrimental. In most of the cases the optimal limit for maximum UTS and elongation is same. As the Y content increases in Mg–6Al and AZ91 alloy, the UTS and elongation increases, but after reaching a maximum corresponding to 0.9 and 1.2 wt% of Y, respectively both start decreasing [16, 17]. The optimal limit

corresponds to the minimum grain size and addition of Y beyond this limit leads to grain coarsening. So, the common interpretation regarding the relation between grain size and mechanical properties holds good. The deterioration in properties after Y exceeds a certain limit is ascribed to the following- (i) preferential growth of large secondary phase in the form of large blocks (ii) the coarse intermetallic phase particles result in easy formation of microcracks and becomes a centre for flaws propagation and (iii) large amount of secondary phase leads to reduction in homogeneity of the microstructure [17]. The addition of Y along with a suitable element in Mg– Al alloys improves the RT properties more significantly. Table 5 shows the effect of binary addition in Mg–6Al alloy. Although the total RE content is 1.8 % in either of the cases, yet the alloy containing a combination of Y and Nd shows higher UTS and elongation at RT. The probable reason for this anomalous behavior may be attributed to the fact that Al2Y phase containing 1.8 wt% Y will be coarser as compared to Al2Y and Al2Nd phase containing 1.2 wt% Y and 0.6 wt% Nd respectively. The later alloy containing harder intermetallics provide a barrier for dislocations to slip and there will be more amount of harder secondary phases through which stress transfer shall take place. 3.2 Mg–Zn Alloys Yttrium plays an important role in improving the RT properties of Mg–Zn alloys as shown in Table 6. It can be deduced that as the Y content increases, both the YS and UTS increase, whereas, the elongation decreases contrary to what is reported that with increasing Y content elongation increases when Zn content is high [40, 50, 51]. In case of Mg–4Zn–0.8Y and Mg–8Zn–1.6Y, the Zn/Y ratio is same but the YS and UTS of the former alloy is higher than that of the later alloy. Also, in case of Mg–4Zn–1.6Y and Mg–8Zn–1.6Y although the Y content is same but the Zn/Y

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Table 4 Room temperature mechanical properties of Y added Mg–Al alloys Alloy

Y (wt %)

UTS (MPa)

Elongation (%)

Reference

AM50

0.0

155

3.6

[32]

AM50

0.3

178

3.8

[32]

Mg–6Al

0.0

190

4.4

[16]

Mg–6Al

0.6

208

6.2

[16]

AZ91

0.0

193

1.0

[60]

AZ91

0.9

220

1.5

[60]

Table 5 Effect of Y addition in mono and binary form in Mg–6Al alloy Alloy

Composition (wt%)

Intermetallic compounds

Grain size (lm)

UTS (MPa)

Mg–6Al

RE = 1.8 (Y = 1.8)

Al2Y

206.8

216

Mg–6Al

RE = 1.8 (Y = 1.2, Nd = 0.6)

Al2Y, Al2Nd

45

253

Elongation (%) 6.2 13

Reference [16] [31]

Table 6 Dependence of room temperature properties of Y added Mg-Zn-Y alloys [42] Zn (wt%)

Y (wt%)

Mg (wt%)

YS (MPa)

UTS (MPa)

Elongation (%)

Phases

3

0.6

Rem.

122

226

30.2

a-Mg ? I-phase

4

0.6

Rem.

132

240

29.5

a-Mg ? I-phase

4

0.8

Rem.

142

245

27.6

a-Mg ? I-phase

8

1.6

Rem.

169

270

26.9

a-Mg ? I-phase

4

1.6

Rem.

162

252

21.7

a-Mg ? I-phase ? W-phase

ratio of the former alloy is much lower than that of the later alloy and so is the YS and UTS. The anomalous variation in properties shows neither Zn/Y ratio nor the Y content can be used accurately to predict the mechanical behaviour of Mg–Zn–Y alloy system. This is because the phases present in these alloys play a vital role in deciding their mechanical behavior, but the formation of phases are neither dictated by the Zn/Y ratio nor the absolute Y content. Mg–4Zn–1.6Y which shows a reduction in YS and UTS in spite of the presence of 1.6 % Y contains W-phase along with a-Mg and I-phase whereas the other alloys showing higher values for YS and UTS contain only a-Mg and I-phase. When an alloy contains quasi-crystals (I-phase) as the second phase, it is stable against the coarsening at high temperature due to low interfacial energy [39] and as a result, no debonding or nanoscale defects are observed in Mg–Zn–Y alloys which are responsible for initiation of failure [37]. Table 6 also reveals that W-phase is detrimental for ductility as the elongation in Mg–4Zn–1.6Y is 21.7 % whereas, in Mg–8Zn–1.6Y it is 26.9 %. When the Zn/Y ratio is so small that only W-phase exists as a secondary phase, then it provides strength to the respective alloy. Figure 4 shows the strengthening effect of W-phase when the Zn/Y ratio is less than 1 in Mg–Zn–Y– Zr alloys. Up to a volume of 1.4 % of W-phase, the

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strength increases and after reaching a maximum, the strength starts deteriorating. The effect of Y content on the properties of Mg–Zn alloys is not always predictable. When Y is added in Mg–5Zn–2Al alloy, up to 0.58 % there is a reduction in UTS, beyond which the UTS starts increasing reaching maximum at 1.78 % [52]. Up to 0.58 % Y the properties deteriorate because of the grain coarsening which may occur as there are a very few Y containing phase and that too distributed within the grains. Moreover, Al atoms start forming Al–Y intermetallic, as a result, there is depletion of Al atoms and increase in Zn atoms solubility in Mg matrix. Therefore, the Mg51Zn20 eutectic phase, which was initially distributed in the grain boundary area as eutectic lumps are now sparsely distributed [52]. There are instances where Y addition causes grain refinement yet the mechanical properties deteriorate. The variation of grain size and corresponding UTS for ZCK 630 with Y addition is shown in Fig. 5. It can be noticed that in spite of Y addition, at none of the compositions the UTS is greater than that of the parent alloy. This abrupt behavior is attributed to the fact that in spite of grain refinement, there is preferable formation of W-phase as compared to I-phase [53] and the former phase easily breaks down which results in reduction in UTS.


Fig. 4 Relationship between the W-phase volume and the mechanical properties for Mg–Zn–Y–Zr alloys with Zn/Y ratio less than 1 [41]

Fig. 5 Variation in UTS of Y added ZCK630 alloys [53]

4 High Temperature Properties In this section the discussion is mainly focused on Mg–Al alloys because of the immense potential possessed by these alloys in automobile and aerospace industries. High Al content leads to higher RT strength, but low microstructure stability at HT [54]. Therefore, majority of work done in Mg–Al alloy system focuses on HT properties since the RT properties lie within the acceptable limits. In Mg–Al alloys, Mg17Al12 phase is responsible for the strengthening effect which mainly resides at the grain boundaries. Since the melting point of Mg17Al12 is 437 °C, its thermal stability is relatively poor which leads to grain boundary weakening at elevated temperature and therefore, it is unable to resist the grain boundary sliding. Also, the rapid diffusion rate of Al at HT results in deterioration of creep properties [55]. Therefore, Mg–Al alloys with low Al content can be used for HT applications as they contain very few Mg17Al12 precipitates. AE42 alloy (Mg–4Al–2RE–0.3Mn) is one of

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the alloys that were specifically developed for high temperature applications. But, the creep properties deteriorate rapidly at HT. Rzychon et al. [56] have reported that as the temperature exceeds 150 °C, there is a partial decomposition of Al11RE3 into Al2RE which leads to release of Al that forms b-Mg17Al12 phase which is detrimental to creep properties. Discontinuous precipitation of Mg17Al12 is another major reason for poor creep properties of Mg–Al alloys. It has been discussed earlier that discontinuous precipitates are a lamellar combination of Mg17Al12 and Al rich a-Mg phase at the grain boundaries and Dargusch et al. [57] have reported that grain boundary sliding is more in lamellar structure as it permits more interfaces. Therefore, in order to obtain better creep properties, there should be not only a reduction in the amount of Mg17Al12 phase, but also its discontinuous mode of precipitation should be suppressed. Two competing trends are generally observed upon Y addition in the traditional Mg–Al alloy systems; first, formation of Al-Y intermetallic phase and second, dissolution of Mg17Al12 secondary phase. The overall creep resistance of Mg–Al alloys is the result of a compromise between the formation and dissolution of the eutectic b-Mg17Al12 phase, formation of thermally stable Al2Y intermetallic compound and solid solution hardening effects of Al in the Mg matrix. In most of the cases, the variation of RT and HT properties are found to follow the same behavior with varying Y content. Initially the properties improve via Y addition and after reaching the maximum, it starts deteriorating in spite of increasing Y content. In fact, the peak point for RT as well as HT properties is also the same being 0.9 wt% Y for AZ91 and 1.2 wt% Y for Mg–6Al alloy as shown in Fig. 6. Small grains are detrimental to creep properties because the microstructure with smaller grain size has a large grain boundary area. The grain boundaries are quasi-viscoelastic in nature and hence start to flow at higher temperature thereby lowering the thermal stability. Further, larger the grain boundary area, more is the grain boundary sliding which is again detrimental to creep. There are various mechanisms that play a collective role in deciding the creep behavior. Wang et al. [17] have reported that following mechanisms collectively decide the creep behavior of AZ91 alloy: (i) precipitation hardening (ii) solution strengthening (iii) changes in grain size in matrix grains and of the size of the precipitates, and (iv) dislocation creep with dispersion strengthening. In Mg–6Al, the initial secondary phase without Y addition is Mg17Al12. However, upon Y addition Al2Y intermetallic is formed whose melting point is 1,485 °C [9] which gets segregated at grain boundaries and reduces the volume fraction of the Mg17Al12 secondary phase, thereby, providing more thermal stability to the grain boundary.

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Fig. 6 The effect of Y addition on a UTS, b elongation at RT and HT in Mg–6Al alloy [13]

AZ91 significantly improves the tensile strength at elevated temperature. One of the reasons attributed to this anomalous behavior is that the addition of Ca along with Y benefits Mg17Al12 as the Mg17Al12 phase containing Ca possesses higher thermal stability than the Mg17Al12 phase without Ca [59].

5 Summary

Fig. 7 Variation in volume of Mg17Al12 and Al2Y phases in Y added AZ80 alloy [21]

Figure 7 shows the variation of Mg17Al12 and Al2Y phases with increasing Y content in AZ80 alloy. As the Y content increases from 0 to 2 %, the volume of Mg17Al12 phase decreases from 11 to 5.5 % whereas the Al2Y increases from 0 to 2.2 %. The observed variations in the volume fractions of the phases can be attributed to the fact that Y consumes some of the Al by forming the Al2Y intermetallic compound, therefore less amount of Al is available to combine with the Mg matrix to form b-Mg17Al12. Also, due to strong ionic and covalent bond interactions and the calculated cohesive energy and density of state (DOS) show that Al2Y has stronger structural stability than Mg17Al12 [58]. When Y is added along with a suitable element, the creep resistance increases more effectively as the effect of harder secondary phases formed add up and leads to substantial reduction in Mg17Al12 and hence more effective pinning of grain boundaries occur. AZ91 with 0.9 wt% Y shows UTS, YS and elongation as 184, 103 MPa and 9 % respectively at 150 °C whereas, with 0.6 wt% Y along with 0.5 wt% Sb the respective values turn out to be 192, 111 MPa and 13 %. The combined addition of Y and Ca in

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The addition of Y in Mg alloys exerts a positive effect on microstructure and mechanical properties. Y readily acts as a site for heterogeneous nucleation for a-Mg phase which makes it a good grain refiner. The addition of Y in Mg alloys also leads to significant increase in UTS as well as ductility at RT and HT. Y on one hand forms intermetallic with Al in case of Mg–Al alloys which have a very high thermal stability and pins up the grain boundaries. On the other hand, it forms ternary phases in Mg–Zn alloys which influence the properties in respective alloy systems. Zn/Y ratio decides the secondary phase that will be formed. More the volume fraction of I-phase higher is the UTS and elongation. But in all the cases, whether it is microstructure, RT or HT properties there is an optimum level of Y addition that leads to improvement in Mg alloy systems. Addition of Y beyond this critical limit is detrimental and leads to grain coarsening which is detrimental for mechanical properties. In order to obtain superior properties addition of a suitable element along with Y can be done, which leads to the formation of one additional intermetallic phase which helps in increasing the efficiency of grain refinement and therefore properties are improved to a higher degree. Acknowledgments The authors wish to thank the Ministry of Human Resource Development, Government of India for the fellowship to the first author and members of Light Meatl Alloys and Composites group for their help.


References 1. Wu R Z, Qu Z K, and Zhang M L, Rev Adv Mater Sci 24 (2010) 35. 2. Hassan S F, Ho K F, and Gupta M, Mater Lett 6 (2004) 2143. 3. Shin-ichi Y, Hyang-Yeon K, Hisamichi K, Akihisa I, and Yoshiaki A, J Alloys Compd 12 (2002) 239. 4. Gang S, Keyna O, Brian C, John W, and Richard H, Mater Sci Eng A 5 (2001) 612. 5. Lu L, Lai M O, and Hoe M L, Nat Mater 4 (1998) 551. 6. Chi H Z, Chen C P, Chen L X, and Wang D Q, J Alloys Compd 10 (2003) 312. 7. Zhao D X, Liu Y H, Shen D Z, Zhang Y J, Lu Y M, and Fan X W, J Cryst Growth 2 (2003) 163. 8. Wang Q D, Lu Y Z, Zeng X Q, Ding W J, Zhu Y P, Li Q H, and Lan J, Mater Sci Eng A 11 (1999) 109. 9. Socjusz-Podosek M, and Lity´nska L, Mater Chem Phys 80 (2003) 472. 10. Penga Q, Mengb J, Lic Y, Huanga Y, and Hort N, Mater Sci Eng A, 528 (2011) 2106. 11. Gao L, Chen R, and Han E, J Alloys Compd 481 (2009) 379. 12. Luo A A, Int Mater Rev 49 (2003) 13. 13. Blum W, Watzinger B, and Zhang P, Adv Eng Mater 2 (2000) 349. 14. Wu G, Fan Y, Gao H, Zhai C, and Zhu Y P, Mater Sci Eng A 408 (2005) 255. 15. Terada Y, Ishimatsu N, Sato R, Sato T, and Ohori K, Mater Sci Forum 419 (2003) 181. 16. Wenliang R, Quan’an L, Jianhong L, Kejie L, and Xingyuan Z, China Foundry 7 (2010) 362. 17. Wang S R, Guo P, Yang L Y, and Wang Y, J Mater Eng Perform 18 (2009) 137. 18. Zhu S M, Gibson M A, Nie J F, Easton M A, and Abbott T B, Scr Mater 58 (2008) 477. 19. B. Amir Esgandari B, Mehrjoo H, Nami B, and Miresmaeili S M, Mater Sci Eng A 528 (2011) 5018. 20. Chu-ming L, Xiu-rong Z, and Hai-tao Z, Phase Diagrams for Magnesium Alloys [M], Central South University Press, Changsha (2006). 21. Nayyeri G, and Mahmudi R, Mater Sci Eng A 527 (2010) 669. 22. Xiao-feng H, Qu-dong W, Xiao-qin Z, Guang-yin Y, Yan-ping Z, and Wei-jiang D, J Chin Rare Earth Soc 22 (2004) 493. 23. Fu-sheng P, Mei-bao C, Jing-feng W, Jian P, and Ai-tao T, Trans Nonferrous Met Soc China 18 (2008) s1. 24. Luo Z P, Song D Y, and Zhang S Q, J Alloys Compd 230 (1995) 109. 25. Lee Y C, Dahle A K, and St John D H, Metall Mater Trans A 31 (2000) 2895. 26. Hongmei L, Yungui C, Yongbai T, Deming H, Min Z, and Yiguo L, J Sichuan Univ 38 2006 90. 27. Dahle A K, Lee Y C, Nave M D, Schaffer P L, and St John D, J Light Met 1 (2001) 61. 28. Polmear I, Light Alloys From Traditional Alloys to Nanocrystals, Butterworth- Heinemann, Boston (2006). 29. Shengfa L, Huiyuan W, Liugen K, Shangyu H, and Ping X, Chin J Nonferrous Met 16 (2006) 464.

339 30. Zhenhua C Wrought Magnesium Alloy, Chemical Industry Press, Beijing, (2007) 25. 31. Jun C, Quanan L, Jianghong L, Xiaofeng L, Kejie L, and Xingyuan Z, China Foundry 6 (2008) 124. 32. Mingxing W, Hong Z, and Wang L, J Rare Earth 25 (2007) 233. 33. Xi-ya F, Dan-qing Y, Bin W, Wen-hai L, and Wei L, Trans Nonferrous Met Soc China 16 (2006) 1053. 34. Gang C, Xiao-dong P, Pei-geng F, Wei-dong X, Qun-yi W, Hong M A, and Yan Y, Trans Nonferrous Met Soc China 21 (2011) 725. 35. Padezhnova E M, Mel’nik E V, Miliyevskiy R A, Dobatkina T V, and Kinzhibao V V, Russ Metall 4 (1982) 185. 36. Luo Z P, and Zhang S Q, J Mater Sci Lett 12 (1993) 1490. 37. Janot C, Quasicrystals, Clarendon Press, Oxford (1994). 38. Pierce F S, Poon S J, and Guo Q, Science 261 (1993) 737. 39. Dubois J M, Plaindoux P, Berlin-Ferre E, Tamura N, and Sordelet D J, Proceedings of the Sixth International Conference on Quasicrystals, World Scientific, Singapore (1997). 40. Xu D K, Liu L, Xu Y B, and Han E H, J Alloys Compd 426 (2006) 155. 41. Xua D K, Tang W N, Liu L, Xu Y B, and Han E H, J Alloys Compd 461 (2008) 248. 42. Lee Y J, Kim D H, Lim H K, and Kim D H, Mater Lett 59 (2005) 3801. 43. Zhang E, He W, Du H, and Yang K, Mater Sci Eng A 488 (2008) 102. 44. Xu D K, Tang W N, Liu L, Xu Y B, and Han E H, J Alloys Compd 432 (2007) 129. 45. Bhan S, and Lal A, J Phase Equilib 14 (1993) 634. 46. Zhou H T, Zhang Z D, Liu C M, and Wang Q W, Mater Sci Eng A 445–446 (2007) 1. 47. Zhang Y, Zeng X, Liu L, Lu C, Zhou H, Li Q, and Zhu Y, Mater Sci Eng A 373 (2004) 320. 48. Ma C, Liu M, Wu G, Ding W, and Zhu Y, Mater Sci Eng A 349 (2003) 207. 49. Hong-hui Z, Trans Nonferrous Met Soc China 18 (2008) 580. 50. Li Q, Wang Q D, Wang Y X, Zeng X Q, and Ding W J, J Alloys Compd 427 (2007) 115. 51. Xu D K, Liu L, Xu Y B, and Han E H, Mater Sci Eng A 443 (2007) 248. 52. Zou H, Zeng X, Zhai C, and Ding W, Mater Sci Eng A 402 (2005) 142. 53. Yun B, Can-feng F, Hai H, Guo-hong Q I, and Xing-guo Z, Trans Nonferrous Met Soc China 20 (2010) s357. 54. Kondori B, and Mahmudi R, Mater Sci Eng A 527 (2010) 2014. 55. Mahmudi R, Kabirian F, and Nematollahi Z, Mater Des 32 (2011) 2583. 56. Rzychon T, and Kielbus A, JAMME 7 (2006) 149. 57. Dargusch M S, Dunlop G L, and Pettersen K, in: Mordike B L, and Kainer K U (Eds), Magnes Alloys Appl, Frankfurt (1998). 58. Zhi-wei H, Yu-hui Z, Hua H, Yu-hong Z, Xiao-feng N, and Peide H, J Cent South Univ 19 (2012) 1475. 59. Feng W, Yue W, Ping-li M, Bao-yi Y U, and Quan-ying G, Trans Nonferrous Met Soc China 20 (2010) s311. 60. Boby A, Ravikumar K K, Pillai U T S, and Pai B C, Proc Eng 55 (2013) 98.

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