Influence of Yttrium Addition on Plastic Deformation of Magnesium

Abstract

Tensile tests of single crystalline and polycrystalline Mg-Y alloys were carried out at room temperature to investigate the influence of yttrium on activation of <c+a> slip systems and to clarify the relationship between ductility of magnesium and the activation of <c+a> slip systems. Tensile directions of single crystals and polycrystals were parallel to (0001) and their rolling direction, respectively. Mg-(0.6–1.1)at%Y alloy single crystals yielded due to the first order pyramidal <c+a> slip (FPCS). Yield stress and ductility of Mg-(0.6–1.1)at%Y alloy single crystals were higher than those of pure magnesium. Mg-0.9at%Y alloy polycrystals showed higher ductility than pure magnesium. The number of grains where second order pyramidal slips were activated was the largest in those where non-basal slips were activated in pure magnesium, while those where FPCS were activated increased with increasing strain in Mg-0.9at%Y alloy. High ductility of Mg-0.9at%Y alloy would be caused by activation of FPCS due to yttrium addition.

 

This Paper was Originally Published in Japanese in J. Japan Inst. Met. Mater. 81 (2017) 458–466.

1. Introduction

Light weight, high specific strength, magnesium has gained attention for use in transport industries such as automobiles, trains, and aircraft. However, magnesium shows low ductility at room temperature. Here, in von-Mises' criterion¹⁾, five independent slip systems are required to deform a crystal to any arbitrary shape. Since the number of $\{0001 \}{<}11 \bar 20{>}$ basal slips²⁾ (hereafter, BS) is insufficient for the criterion, non-basal slips must be activated, such as $\{10 \bar 11 \}{<}11 \bar 23{>}$ first order pyramidal <c+a> slip (FPCS) or $\{11 \bar 22 \}{<}11 \bar 23{>}$ second order pyramidal <c+a> slip (SPCS). Ando et al.^3–7) has reported that pure magnesium single crystals yielded and deformed due to activation of SPCS in both $[11 \bar 20]$ tensile and [0001] compression tests at room temperature. Therefore, controlling <c+a> slip systems is a key to improving ductility of magnesium.

It has been reported that the addition of yttrium, a rare earth element, improves room temperature ductility of magnesium^8,9). Mineta et al.¹⁰⁾ performed pure shear tests of Mg-0.8at%Y alloy single crystals and reported that critical resolved shear stress (CRSS) for $\{10 \bar 10 \}{<}11 \bar 20{>}$ prismatic slip (PS) was lower than that of pure magnesium. However, activation of PS by itself is insufficient for von-Mises' criterion. Also, Sandlöbes et al.¹¹⁾ has reported that Mg-3.0mass%Y alloy polycrystals at room temperature had five times higher ductility than pure magnesium and strength as same as pure magnesium. In Ref. 11), <a> dislocations induced by activation of BS and large number of <c+a> dislocations were observed by TEM after tensile tests. Therefore, such high ductility would be caused by activation of <c+a> dislocations through addition of yttrium. However, details of <c+a> slip systems and their influence by yttrium addition have yet to be elucidated. In this study, tensile tests of Mg-Y alloy single crystals and polycrystals were carried out to investigate influence of yttrium on <c+a> slip systems and the relationship between ductility of magnesium and activation of <c+a> slip systems.

2. Experimental Procedures

2.1 Tensile tests of single crystals

Mg-0.9at%Y and Mg-1.1at%Y alloy ingots were produced using a high frequency induction furnace. Ingots and seed crystals were installed in a graphite mold and then sheet shaped Mg-(0.6, 0.9, 1.1 and 1.3)at%Y single crystals were grown by the Bridgeman method. The composition of each single crystal was analyzed by ICP emission spectroscopy. Single crystals were crystallographically analyzed by the X-ray back reflection Laue method, and were then cut with a non-distortion cutting machine using a stainless wire and nitric acid. The surfaces of specimens were chemically polished using a polishing cloth which was a soaked in a 20 ml C₂H₅OH + 7 ml H₂O₂ + 5 ml HNO₃ solution. The polished specimens were then annealed in a thermal cyclic to remove the strain induced from cutting and polishing. For each cycle, the polished specimens were annealed between 673 and 723 K; they were held for 3.6 ks at both temperatures with the rate of change between temperatures 6.9 × 10⁻³ K/s. Four types of Mg-Y alloy single crystal specimens with different crystal orientations were prepared for tensile tests, as shown in Fig. 1. Four different orientations were chosen in order to consider the influence of Schmidt factor on non-basal slips. The Schmidt factor for SPCS is highest when the tensile direction is parallel to ${<}11 \bar 20{>}$ and lowest when the tensile direction is parallel to $[10 \bar 10]$, but for PS, highest when the tensile direction inclines by 15 degrees from $[11 \bar 20]$ toward $[10 \bar 10]$. For SPCS, the A, A13 and G specimens had identical surface plane parallel to basal planes but different tensile directions, which were for A, A13 and G respectively $[11 \bar 20]$, 13 degrees tilted from $[11 \bar 20]$ toward $[10 \bar 10]$ and $[10 \bar 10]$. B specimen had the surface plane parallel to $(10 \bar 10)$ and tensile direction parallel to $[11 \bar 20]$. Two stainless chucks with 10 mm in length, 5 mm in width and 5 mm in thickness were attached to each single crystal specimen using an adhesive for tensile tests, as shown in Fig. 2 (a). Tensile tests were carried out at room temperature with the crosshead speed at 8.3 × 10⁻⁴ mm/s. The value for initial strain rates used in this study were crosshead speed divided by the length of specimen L to produce $\dot \varepsilon $, which were 4.4～8.0 × 10⁻⁵/s. Orientations, compositions, dimensions, initial strain rates $\dot \varepsilon $, yield stresses σ_y and fracture strains ε_f are summarized in Table 1. Specimen names mean its composition, orientation and test numbers.

Fig. 1

Schematic illustrations of A, A13, B and G-specimens used in this study and hcp unit cells are also drawn.

Fig. 2

Schematic illustrations of tensile test specimen for (a) single crystals and (b) polycrystals.

Table 1 Compositions, dimensions, initial strain rates and mechanical properties of Mg-Y alloy single crystals.

Specimen	Composition (at%)	L/mm	W/mm	T/mm	$\dot \varepsilon /{\rm s}^{-1}$	σy/MPa	εf/%
0.6A-1	0.6	18.74	2.66	0.24	4.4 × 10−5	140	6.3
0.6A-2		17.65	2.72	0.24	4.7 × 10−5	141
0.9A13-1	0.9	10.34	2.81	0.10	8.0 × 10−5	221	26.7
0.9A-1		11.94	2.98	0.22	7.0 × 10−5	171	15.7
0.9A-2		10.79	2.94	0.21	7.7 × 10−5	172	10.2
1.1A-1	1.1	12.67	3.14	0.25	6.6 × 10−5	226	5.4
1.1A-2		11.83	3.28	0.17	7.0 × 10−5	230	13.0
1.1A-3		12.31	3.01	0.17	6.7 × 10−5	223	0.8
1.1B-1		13.85	3.03	0.17	6.0 × 10−5	137
1.1B-2		14.17	2.83	0.12	5.9 × 10−5	158	1.4
1.1B-3		12.60	2.94	0.17	6.6 × 10−5	180	2.3
1.3G-1	1.3	15.18	3.14	0.17	5.5 × 10−5	173	6.5
1.3G-2		14.57	3.08	0.17	5.7 × 10−5	183

2.2 Tensile tests of polycrystals

Casted Mg-0.9at%Y alloys were hot-rolled at 523 K. The rolling reduction was approximately 10% per one pass, and total reduction was 75%. Pure magnesium was also rolled under the same condition to use for comparison. To obtain grains with grain sizes of approximately 60 μm, rolled Mg-0.9at%Y alloy sheets and rolled pure magnesium sheets were annealed at 773 K for 480 s and at 673 K for 360 s. Annealed sheets were cut into tensile specimens, as shown in Fig. 2 (b). Tensile specimens were mechanically polished using emery papers (#400–4000) and a 1 μm DP spray (Struers) with a MD-Nap (Struers) and were soaked into the same chemical solution used for single crystals to obtain mirror surfaces. Crystal orientations in normal direction (ND) plane of specimens were analyzed by electron backscattered diffraction (EBSD) method. Tensile tests were carried out at room temperature. The crosshead speed and $\dot \varepsilon $ were 1.7 × 10⁻² mm/s and 8.3 × 10⁻⁴/s, respectively. Composition, mean grain sizes and dimensions of the specimens are summarized in Table 2. Slip lines and twins were observed with a Nomarsky type optical microscope.

Table 2 Compositions, grain sizes, dimensions and mechanical properties of pure magnesium and Mg-Y alloy polycrystals.

Specimen	Composition	Grain size/μm	GL/mm	W/mm	T/mm	σy/MPa	εf/%
M-1	Pure magnesium	60.7	20.0	3.83	2.24	80	4.0
M-2		56.9	20.0	3.97	1.83	68
Y-1	Mg-0.9at%Y	59.2	20.0	4.00	2.49	97	20.0
Y-2		59.2	20.1	3.94	2.48	100

3. Results

3.1 Deformation behavior of single crystals

Figure 3 shows typical stress-strain curves of Mg-Y alloy single crystals. Arrows in Fig. 3 indicate yield and fracture points. In this study, yield points are defined as points at which the stress-strain curve deviates from the linear relationship between stress and strain. 0.6A-1, 0.9A13-1 and 1.1A-1 yielded at 140 MPa, 221 MPa and 226 MPa and deformed without work hardening. Here, the composition and tensile direction of 1.1B-3 and 1.1A-1 were nearly the same. However, yield stress of 1.1B-3 was 180 MPa, much lower than that of 1.1A-1. Also, the flow stress of 1.1B-3 suddenly dropped after yield, and the elongation was much lower than that of 1.1A-1. 1.3G-1 yielded at 173 MPa and showed small work hardening. Thus, the deformation mechanism of 1.3G-1 would differ from that of other specimens. In $[11 \bar 20]$ tensile tests, the yield stress and fracture strain of pure magnesium have been reported to be approximately 80 MPa and 3.2%⁵⁾. Therefore, addition of yttrium would increase yield stress and fracture strain.

Fig. 3

Typical stress-strain curves of Mg-Y alloy single crystals.

Figure 4 shows optical micrographs of 0.6A-1, 0.9A13-1, 1.1B-3 and 1.1G-2 after tensile tests. 0.6A-1 shows necking resulting from instability during tensile deformation. Necking is also observed in 0.9A13-1, but the area is wider than that in 0.6A-1. Therefore, the entire deformation of 0.9A13-1 can be considered far greater than that of 0.6A-1. In 1.1B-3, slip lines were not observed on the surface of the whole specimen but were limited to areas around a crack, indicating it deformed locally and broke early in the tensile test. On the other hand, slip lines - but not necking - were observed in the whole specimen of 1.3G-2. Therefore, 1.3G-2 would deform uniformly in the tensile test. Typical slip lines observed on (0001) are shown in Fig. 5. Here, arrows in Fig. 5 indicate initial orientations of specimens. Slip lines were observed on the surface of 0.6A-2 at a strain of 0.3% just after yield, as shown in Fig. 5 (a). That the slip lines would be due to FPCS is known as its direction tilted 60 degrees from $[11 \bar 20]$. Also, the slip lines grow more noticeable with increasing strain, as shown in Fig. 5 (b). 0.6A-2 was expected to deform solely by activation of FPCS from the yield to fracture points because other slip lines were not observed. Similar slip lines were also observed in both 0.9A13-1 and 1.1A-2, as shown in Fig. 5 (c) and (d). Therefore, in the quantity of yttrium content ranging from 0.6 to 1.1at%, Mg-Y alloy deforms solely by activation of FPCS.

Fig. 4

Optical micrographs of (a) 0.6A-1, (b) 0.9A13-1, (c) 1.1B-3 and (d) 1.3G-2 after tensile tests.

Fig. 5

Optical micrographs of (0001) surfaces of (a) 0.6A-2, (b) 0.6A-2, (c) 0.9A13-1 and (d) 1.1A-2.

Figure 6 shows optical micrographs of (a) 1.1B-3 and (b) 1.3G-2 after tensile tests. Slip lines were clearly observed on $(1 \bar 100)$ in 1.1B-3. The slip lines would be PS because they are perpendicular to $[11 \bar 20]$. On the other hand, in 1.3G-2, slip lines tilted 30 and 90 degrees from $[1 \bar 100]$ were respectively observed on (0001) and $(11 \bar 20)$, as shown in Fig. 6 (b). Thus, the slip lines are geographically analyzed to be PS as well in 1.1B-3. Therefore, when the quantity of yttrium content exceeds 1.1at%, Mg-Y alloys deform due to activation of PS.

Fig. 6

Optical micrographs of (a) 1.1B-3 and (b) 1.3G-2. Slip lines are parallel to [0001], and they were caused by activation of PS.

Figure 7 shows that the relationship between CRSS for non-basal slips and yttrium content of Mg-Y alloy single crystals. FPCS was activated in 0.6A, 0.9A, 0.9A13 and 1.1A, as shown in Figs. 4 and 5. CRSS for FPCS was calculated using Schmidt factors and yield stresses of 0.6A, 0.9A, 0.9A13 and 1.1A respectively, with resulting values approximately 57 MPa, 70 MPa, 75 MPa and 90 MPa. We calculated CRSS for SPCS and FPCS using tensile data of magnesium single crystals reported by Ando et al.^3,6), and they were approximately 40 MPa and 53 MPa. Since SPCS was not activated in 0.6A, 0.9A, 0.9A13 and 1.1A, CRSS for SPCS would significantly increase by yttrium addition. It was also found that CRSS for FPCS increased with increasing yttrium content. In addition, PS was activated in 1.1B and 1.3G, and the CRSS were respectively calculated to be approximately 70 MPa and 80 MPa. The nearly identical value indicates that CRSS for PS would be independent of the quantity of yttrium addition. Therefore, Mg-Y alloys deformed due to activation of PS because CRSS for FPCS was larger than that for PS when the addition of yttrium is more than 1.1at%.

Fig. 7

Relationship between CRSS for non-basal slips and yttrium content of Mg-Y alloy single crystals.

Figure 8 shows the relationship between tensile elongation and yttrium content of Mg-Y alloy single crystals. Specimens in which FPCS was activated showed higher ductility than those in which SPCS or PS was activated. Specimens with 0.9at% yttrium showed the highest ductility in this study, indicating that FPCS plays a role of increase in ductility.

Fig. 8

Relationship between tensile elongation and yttrium content of Mg-Y alloy single crystals.

3.2 Deformation behavior of polycrystals

Figure 9 shows typical stress-strain curves of magnesium polycrystals. M-1, pure magnesium polycrystals, yielded at 80 MPa and then showed rapid work hardening. On the other hand, Y-1, Mg-0.9at%Y alloy polycrystals, yielded at 97 MPa and showed gradual work hardening behavior. As a result, Y-1 showed five times higher ductility than M-1, and the result is similar to that reported by Sandlöbes et al.¹¹⁾

Fig. 9

Typical stress-strain curves of M-1 and Y-1.

Figure 10 shows optical micrographs (a) and (b), inverse pole figure (IPF) maps (c) and (d), and (0002) pole figures (e) and (f) of M-1 and Y-1 specimens before tensile tests. M-1 and Y-1 had equiaxed grains and showed basal texture with c-axes parallel to ND. However, Y-1 had more random orientation than M-1, as shown in Fig. 10 (e) and (f). The result is in agreement with that reported by Suzuki et al.¹²⁾ Thus, yttrium addition makes basal texture weaker.

Fig. 10

Optical micrographs, IPF maps and (0002) pole figures of M-1 (a), (c) and (e) and Y-1 (b), (d) and (f) before tensile tests.

Figure 11 shows optical micrographs of (a) M-1 at a strain of 4.0% and (b) Y-1 at a strain of 5.2% after tensile tests. The observation areas are as same as those shown in Fig. 10 (a) and (b). Arrows in Fig. 11 (a) and (b) indicate twins. Slip lines were observed within a half number of grains in M-1, although they were not clearly seen in Fig. 11 (a). The enlargement is shown in Fig. 12 (a) and (b) later. Twins were also observed within many grains in M-1. The number of twins in Y-1 was smaller than that in M-1. Slip lines were observed within approximately 97% of the number of grains, shown in Fig. 11 (b). Reasons why the microstructure was much different in of M-1 and Y-1 after tensile tests are described below. In M-1, strong basal texture formed, and the basal plane was parallel to the tensile direction. As a result, Schmidt factor for BS was close to zero and BS was barely activated. On the other hand, because the basal texture weakened due to yttrium addition, BS, the main slip system, was more easily activated. In order to distinguish activated slip systems, slip lines were geographically analyzed. Figure 12 shows optical micrographs of slip lines observed in different areas at a strain of 4.0% in M-1. Slip lines caused by BS were observed within grains shown in Fig. 12 (a) and (b). However, slip lines caused by FPCS and SPCS were observed within several grains and slip lines caused by PS were never observed. Figure 13 shows optical micrographs of slip lines at ε = 1.7%: (a), (b) and (c) and 5.2%: (d), (e) and (f) in Y-1. Slip lines were observed in almost all grains - not only BS but also FPCS, SPCS and PS were activated in Y-1. Non-basal slip lines were seemed to connect with BS lines within neighboring grains at ε = 1.7%, as shown in Fig. 13 (a), (b) and (c). On the other hand, at ε = 5.2%, slip lines appear independent as shown in Fig. 13 (d), (e) and (f). Activation of FPCS, SPCS and PS are discussed below.

Fig. 11

Optical micrographs of (a) M-1 when ε was 4.0% and (b) Y-1 when ε was 5.2%.

Fig. 12

Optical micrographs of slip lines caused by basal and non-basal slips: (a) FPCS and (b) SPCS on the surface of M-1 when ε was 4.0%.

Fig. 13

Optical micrographs of basal and non-basal slip lines on the surface of Y-1 when ε was 1.7% (a), (b) and (c) and 5.2% (d), (e) and (f).

4. Discussion

Figure 14 shows the relationship between the frequency of non-basal slips and nominal strain in M-1 and Y-1. Here, the frequency of non-basal slips is defined as the ratio of the number of grains with slip lines to the number of all grains: 190 grains in M-1 and 152 grains in Y-1. In M-1, SPCS was most activated while activation of FPCS and PS were limited. Frequency of FPCS, SPCS and PS were 12%, 10% and 14% at ε = 1.7% in Y-1, respectively. The frequency of PS and FPCS increased with increasing strain, but that of SPCS barely increased. It was therefore found that PS and FPCS were more active than SPCS in Y-1 and that the frequency of non-basal slips in Y-1 was higher than that in M-1. Additionally, the frequency of BS in M-1 and Y-1 were 49% and 97%, respectively.

Fig. 14

Relationship between the frequency of non-basal slips and nominal strain.

Schmidt factors for activated slip systems were calculated based on orientation of grains analyzed by the EBSD method. Resolved shear stresses (RSSs) for non-basal slips were calculated using the calculated Schmidt factor and the loading stress in tensile tests. Figure 15 shows relationships between RSSs for non-basal slips and nominal strain. Dashed lines in Fig. 15 indicate CRSSs for FPCS, SPCS and PS. CRSSs for Mg-0.9at%Y alloy were calculated from tensile data of single crystals in this study. CRSSs for pure magnesium were calculated using data reported by Ando et al.³⁾, as described above. SPCS were activated at ε = 4.0% in M-1 because RSS for SPCS was higher than CRSS in that situation; conversely, FPCS and PS were activated in Y-1, although RSSs for FPCS and PS were lower than CRSS for those at ε = 1.7%. Thus, non-basal slips were activated due to stress concentration at grain boundaries by activation of BS during the initial deformation stage just after yield, as shown in Fig. 13 (a), (b) and (c). However, RSSs for non-basal slips were higher than CRSS for those at ε = 5.2% and non-basal slips would be activated solely, as shown in Fig. 13 (d), (e) and (f).

Fig. 15

Relationship between the resolved shear stress (RSS) on non-basal slips and nominal strain.

In pure magnesium, SPCS are easily activated because CRSS for SPCS is lower than that for FPCS. However, it has been reported that pure magnesium single crystals show rapid work hardening due to activation of SPCS and break at a tensile strain of approximately 3.0%⁵⁾. Therefore, M-1, pure magnesium polycrystals, would show rapid work hardening due to activation of SPCS during tensile tests and break at only several percent strain. On the other hand, it was found that Mg-0.9at%Y single crystals deformed due to activation of only FPCS. While CRSS for FPCS is higher than that for SPCS of pure magnesium, Mg-0.9at%Y alloy single crystals show high ductility even without most any work hardening. Therefore, in Y-1, FPCS may be activated due to the stress concentration at grain boundaries by activation of BS just after yield. With increasing deformation, FPCS may be solely activated because RSS for FPCS surpasses CRSS in that case. That is, FPCS is more highly activated by the addition of yttrium since the CRSS for FPCS becomes lower than that for SPCS. Also, tensile data of 0.9A13 show that activation of FPCS plays a role of increasing ductility. It was therefore found that high ductility of Y-1 causes by activation of FPCS, $\{10 \bar 11\}{<}11 \bar 23{>} $ first order pyramidal <c+a> slip.

5. Conclusions

Tensile tests of Mg-Y alloy single crystals and polycrystals were carried out at room temperature to investigate influence of yttrium on activation of <c+a> slip system and the relationship between ductility of magnesium and activation of <c+a> slip systems. The main results are as follows:

• (1)   Mg-(0.6～1.1)at%Y alloy single crystals yielded due to $\{10 \bar 11 \}{<}11 \bar 23{>}$ first order pyramidal <c+a> slip and deformed without work hardening.
• (2)   Mg-(1.1～1.3)at%Y alloy single crystals yielded due to $\{10 \bar 10 \}{<}11 \bar 20{>}$ prismatic slip and deformed with work hardening.
• (3)   CRSS for $\{10 \bar 11 \}{<}11 \bar 23{>}$ first order pyramidal <c+a> slip increased with increasing yttrium addition.
• (4)   In Mg-0.9at%Y alloy polycrystals, non-basal slips could be activated at lower resolved shear stress than its CRSS due to stress concentration at grain boundaries by activation of basal slips.
• (5)   In Mg-0.9at%Y alloy polycrystals, the frequency of $\{10 \bar 11 \}{<}11 \bar 23{>}$ first order pyramidal <c+a> slip and $\{10 \bar 10 \}{<}11 \bar 20{>}$ prismatic slip increased with increasing strain.
• (6)   High ductility of Mg-0.9at%Y alloy polycrystals was caused by activation of $\{10 \bar 11 \}{<}11 \bar 23{>}$ first order pyramidal <c+a> slip from just after yield to fracture.

Acknowledgements

A portion of the present study was financially supported by “The Light Metal Educational Foundation, Inc.”. The authors are very grateful for the support.

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