The Portevin–Le Chatelier effect in β-phase Mg–14.3Li–0.8Zn alloy

Materials Science & Engineering A 605 (2014) 33–38

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The Portevin–Le Chatelier effect in β-phase Mg–14.3Li–0.8Zn alloy S.K. Wu a,b,n, C. Chien b, C.S. Yang a, H.Y. Bor c a b c

Department of Mechanical Engineering, National Taiwan University, Taipei 106, Taiwan Department of Materials Science and Engineering, National Taiwan University, Taipei 106, Taiwan Metallurgy Section, Materials and Optoelectronics Research Division, Chung-Shang Institute of Science and Technology, Taoyuan 325, Taiwan

art ic l e i nf o

a b s t r a c t

Article history: Received 9 November 2013 Received in revised form 21 January 2014 Accepted 6 March 2014 Available online 17 March 2014

The effects of strain rate, ε_ , and temperature, T, on the occurrence of the Portevin–Le Chatelier (PLC) effect in tensile hot-rolled (HR) and solid-solution treated (SS) β-phase LZ141 magnesium alloy were studied. The HR alloy has an intrinsically higher dislocation density and fewer solute atoms than the SS alloy. This characteristic, in terms of the dynamic strain aging (DSA) mechanism, explains why the PLC effect does not occur in HR alloy for ε_ ranging from 3.33
10  4 s  1 to 6.67
10  2 s  1 but does occur in SS alloy, in which Type B and Type C serrations appear at ε_ ¼ (3.33–6.67)
10  3 s  1 and at ε_ ¼ (3.33– 6.67)
10  4 s  1, respectively. The SS alloy exhibits a negative strain rate sensitivity (SRS) at room temperature. The negative SRS also supports the proposition that the DSA mechanism causes the PLC effect. In the study of the effect of T on the occurrence of the serrated flow, for HR alloy at Tr 0 1C, Type A serrations were observed. In contrast, in the SS alloy, Type C serrations occurred in the curves at 25 1C and 0 1C, and Type B serrations occurred at  25 1C and  50 1C. These results can also be explained by the DSA mechanism. Large serrated stress variations were found in the tensile curves of SS alloy at  25 1C and  50 1C, but no twinning was found near the fractured surface. & 2014 Elsevier B.V. All rights reserved.

Keywords: The Portevin–Le Chatelier effect Mechanical characterization Magnesium alloy Bulk deformation Plasticity

1. Introduction Mg–Li alloys are less dense than magnesium itself, so they have been the subject of much study in the past. Mg–Li alloys that contain more than E11% Li have a single β-phase with a BCC structure, and those that contain E5–11% Li consist of two phases, α þ β, with the α-phase having an HCP structure [1–3]. (The composition of the alloy used in this study is given in wt%.) The Portevin–Le Chatelier (PLC) effect, or the so-called “serrated flow”, has been reported in α-phase magnesium alloys. In the 1960s and 1970s, Mg–0.5% Th alloy was reported to have an anomalous yielding effect [4] and serrated flow in the temperature range 240–285 1C [5]. Serrated flow was also observed in Mg–10% Ag alloy at 53–124 1C in solid-solution treated specimens [6,7]. In the past decade, more α-phase magnesium alloys have been found to exhibit the PLC effect under certain conditions, such as AZ91 (Mg–9% Al–1% Zn) [8], WE54 (Mg–(5.0–5.5)% Y–(1.5–2.0)% Nd– (1.5–2.0)% RE–0.4% Zr) [9], and LA41 (Mg–4.32% Li–0.97% Al) [10]. For Mg–Al–Zn (AZ) alloys, Corby et al. reported that only AZ91 alloy exhibits serrated flow at room temperature [8]. For WE54

n Corresponding author at: Department of Materials Science and Engineering, National Taiwan University, No. 1, Sec. 4, Roosevelt Road, Taipei 106, Taiwan. E-mail address: [email protected] (S.K. Wu).

http://dx.doi.org/10.1016/j.msea.2014.03.026 0921-5093/& 2014 Elsevier B.V. All rights reserved.

alloy, Zhu and Nie [9] observed serrated flow in the temperature range 150–225 1C and attributed it to the dynamic strain aging (DSA) effect [11]. For LA41 alloy containing o5% Li, Wang et al. [10] found that serrated flow is apparent throughout the tensile deformation and that this alloy exhibits abnormal strain rate sensitivity (SRS), with SRS being positive in the strain rate (ε_ ) range 1.33
10  4–6.67
10  4 s  1, and negative in the ε_ range 6.67
10  4–1.33
10  2 s  1. The variation in SRS is thought to result from competition between the DSA of solute atoms and the shearing of precipitation by dislocations. Many studies have been focused on Mg–Li β-phase alloys that contain more than 12% Li because of their intrinsic ultralight property. Clark and Sturkey reported that, for Mg–19.5Li–18.7Zn alloy aged at 21–150 1C, a transition structure θ0 (MgLi2Zn) phase is generated upon precipitation of the stable LiZn phase, which is observed in the Debye– Scherrer X-ray pattern [12]. Levinson and McPherson studied Mg– 12.3Li–1.0Al alloy and found that θ (MgLi2Al) phase is not an equilibrium phase in the temperature range 100–400 1C [13]. Takuda et al. studied tensile Mg–12Li–1Zn alloy and found it is particularly sensitive to ε_ and has sufficiently high ductility at low ε_ [14]. Song et al. reported that, for Mg–12Li–0.03Be–xAl (x¼1, 3) alloys, the θ phase is formed at room temperature in the casts, and its formation is accelerated if the Al content is increased [15]. Liu et al. studied equal channel angular processed (ECAP) Mg–14Li–1Al (LA141) alloy and found that the grains of the β-phase matrix are substantially more

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S.K. Wu et al. / Materials Science & Engineering A 605 (2014) 33–38

refined [16]. Wu et al. studied the mechanical properties of cold-rolled and/or solid-solution treated β-phase Mg–14.3Li–0.8Zn (LZ141) alloy and concluded that the duplex strengthening effect of solid-solution and subsequent severe cold-rolling significantly improves the mechanical properties [17]. Studies of β-phase Mg–Li alloys tend to focus on aspects such as their microstructural observation, the θ/θ0 phase transition, and their mechanical properties. Although the PLC effect has been observed in iron (BCC structure) containing small concentrations of carbon [18], the PLC effect in tensile deformed β-phase Mg–Li alloys has not yet been reported. In this study, the PLC effect was exhibited in the tensile stress–strain curves of β-phase Mg–14.3Li– 0.8Zn (LZ141) alloy. The effect of ε_ and temperature on the occurrence of the PLC effect and the reasons for this phenomenon are discussed.

2. Experimental procedures LZ141 alloy was prepared from the pure raw materials of magnesium (purity 99.95%), lithium (purity 99.9%), and zinc (purity 99.99%). These raw materials were induction-melted and protected by argon gas, cast in a steel mold, and homogenized at 350 1C for 12 h. The optical micrograph (OM), X-ray diffraction (XRD) pattern and scanning electron microscope (SEM) image of as-homogenized LZ141 alloy were shown in our previous work [17]. The precise chemical composition of as-homogenized LZ141 alloy was determined to be Mg–14.3Li–

0.8Zn, with trace amounts of Al, Mn and Fe of less than 0.02%, using an inductively coupled plasma-optical emission spectrometer (ICP-OES). The homogenized ingot was sliced into plates of 30 mm thickness. The plates were then hot-rolled (abbreviated as HR) at 200 1C to plates thickness of 3 mm. The specimens for solid solution treatment (abbreviated as SS) were HR alloy heated to 350 1C for 20 min and then water-quenched. The specimens for microstructural examination were prepared using the standard metallographic procedure with an etching solution of 1 ml 2,4,6-trinitrophenol, 1 ml water, and 7 ml ethanol. Microstructural observations were performed using a Nikon optical microscope. The ASTM E112-88 standard was used to calculate the average grain size [19]. The specimens for tensile tests were produced according to ASTM E8/E8M-13a [20] with a gauge length of 25.0 mm and width of 6.0 mm. Tensile tests were performed using a Shimadzu testing machine (AG-IS 50 kN, Japan) with a stroke control and a strain rate ε_ of 6.67
10  5–6.67
10  2 s  1. XRD measurements of the crystal structure were performed using a Rigaku diffractometer (Rigaku TTR AXIII, Japan) with a Cu Kα X-ray tube operated at 50 kV voltage and 300 mA current at a step width of 0.021 and a measurement time of 0.3 s per step.

3. Results and discussion 3.1. Microstructure observation and XRD measurement Fig. 1(a) and (b) shows OM images of HR and SS LZ141 alloys on the rolling plane, respectively. Fig. 1(a) demonstrates that the

(200)

(211)

80000

Intensity

HR LZ141 β phase α phase

(110)

100000

60000

200μm

(1012)

(1010)

20000

(220)

40000

0 20

30

40

50

60

70

80

(110)

2θ

50000

LZ141 SS β phase

30000

(220)

20000

(220)

10000

(211)

intensity(cps)

40000

0

200μm Fig. 1. Optical micrographs of (a) hot-rolled (HR) and (b) solid-solution treated (SS) LZ141 alloys. The arrow shown in (a) is the rolling direction.

20

30

40

50

60

70

80

2θ Fig. 2. X-ray diffraction (XRD) patterns of (a) hot-rolled (HR) and (b) solid-solution treated (SS) LZ141 alloys.

S.K. Wu et al. / Materials Science & Engineering A 605 (2014) 33–38

grains were elongated along the rolling direction and that the residual plastic deformation in the grains was retained. Fig. 1 (b) shows an equiaxial microstructure with an average grain size of E150 μm. Fig. 2(a) and (b) shows the XRD patterns of Fig. 1(a) and (b), respectively. In Fig. 2(a), it is notable that the HR alloy was mainly β-phase with small α-phase peaks. According to SEM observation of as-homogenized LZ141 alloy [17], the tiny αphase particles exist mostly at the grain boundaries and rarely in the grains. No α-phase peaks can be observed in Fig. 2(b), which means α-phase particles were dissolved in the β-phase of the SS LZ141 alloy. No θ0 -phase peak was found in Fig. 2(a) and (b).

3.2. The effect of strain rate on the tensile stress–strain curve Fig. 3(a) and (b) shows the tensile stress–strain curves for HR and SS LZ141 alloys, respectively, at room temperature with the strain rate ε_ of 6.67
10  5–6.67
10  2 s  1. Fig. 3(a) shows no PLC effect in the stress–strain curves for all strain rates for the HR alloy. Meanwhile, in Fig. 3(b), the PLC effect can be seen in the curves of the SS alloy at some strain rates. This characteristic comes from the fact that the α-phase particles in the HR alloy are dissolved in the β-matrix of the SS alloy, as shown in Fig. 2. These dissolved atoms interact dynamically with the dislocations to generate the serrated flow, as the DSA mechanism first proposed by Cottrell in 1953 [11]. In order to clearly observe the appearance of the serrated flow, the stress–strain curves in Fig. 3(b) were enlarged in the strain range 0–0.20, as shown in Fig. 3(c). In that figure, the PLC effect is visible in the stress–strain curves of SS LZ141 alloy under the

140

100 80

6

3.33x10 6.67x10 3.33x10 6.67x10 3.33x10 6.67x10

s s s s s s

4 3

40

1

2

20

80

0

6.67x10 3.33x10 6.67x10 3.33x10 6.67x10 3.33x10 6.67x10

s s s s s s s

1

4

5

60 40

3

20

0

2

0 0

20

40

60

80

0

20

40

Strain ε (%)

120

6.67x10 3.33x10 6.67x10 3.33x10

s 4 s 5 s 6 s

6.67x10 s 3.33x10 s 6.67x10 s

100

LZ141 SS 1. 3.33x10 2. 6.67x10 3. 3.33x10 4. 6.67x10 5. 3.33x10

80

4

5

1

2 3

100

80

85

6

LZ141 SS 0 1 2 3

60

Strain ε (%)

80

0

Flow Stress σ0.2 (MPa)

140

Stress σ (MPa)

0 1 2 3 4 5 6

100

5

60

LZ141 SS

6

120

Stress σ (MPa)

1 2 3 4 5 6

120

Stress σ (MPa)

applied ε_ range 3.33
10  4 (curve 1)–3.33
10  2 s  1 (curve 5). The serration type at ε_ ¼3.33
10  4 s  1 (curve 1) and ε_ ¼6.67
10  4 s  1 (curve 2) is identified as Type C, in which the serrations fall below the general level of the stress–strain curve [21,22]. Both the serration density and the strain range of the PLC effect are higher in curve 2 than in curve 1. The stress–strain curves for ε_ ¼3.33
10  3 s  1 (curve 3) and ε_ ¼6.67
10  3 s  1 (curve 4) exhibit Type B serrations at smaller strain (ɛ o E 6% for curve 3 and ɛ o E12% for curve 4), in which the serrations oscillate about the general level of the stress–strain curve [21,22]. For the strain range 6–8% in curve 3 and that 12–17% in curve 4, the serration type is Type C, with the serrated stress drop and the serration density being less than those in curves 1 and 2. In addition, curves 3 and 4 show that the stress increases slightly as the serration type changes from Type B to Type C. The stress–strain curve for ε_ ¼3.33
10  2 s  1 (curve 5) exhibits insignificant Type B serrations in the strain range 0.03–0.20, as compared with those in curves 3 and 4. Fig. 3 also shows no PLC effect in ε_ ¼6.67
10  5 (curve 0) and ε_ ¼6.67
10  2 (curve 6). Jiang et al. [23] reported that in 2017 aluminum alloy (Al–4% Cu– 0.6% Mg alloy) under a ε_ ranging from 1
10  5 s  1 to 5
10  3 s  1, Type A serrations occur at a larger applied ε_ because the mobility of solute atoms cannot keep up with the fast dislocation motion, so a weak DSA effect causes slight serrations on the tensile curve. Types B and C occur at a lower applied ε_ , in which the pinning process is more significant in dislocations, so an enhanced DSA effect is responsible for the striking serrations on the tensile curve. Fig. 3 (c) demonstrates that no Type A serrations occur; only Types B and C do. This characteristic indicates that, from the viewpoint of Ref.

LZ141 HR

140

35

1 75

70

s s s s s

3

2

65

4

5

60 60 0

5

10

Strain ε (%)

15

20

1E-4

1E-3

0.01 1E-2

0.1 1E-1

log(ε) (s-1)

Fig. 3. Tensile stress–strain curves of (a) hot-rolled (HR) and (b) solid-solution treated (SS) LZ141 alloys tested at room temperature under strain rates ε_ of 6.67
10  5– 6.67
10  2 s  1. (c) The enlarged curves of (b) in the strain range 0–0.20. (d) The plot of the flow stress vs. ln ε_ for curves 1–5 is shown in (c).

S.K. Wu et al. / Materials Science & Engineering A 605 (2014) 33–38

[23], there is no weak DSA effect, but notably, there is an enhanced DSA effect that causes the PLC effect in SS LZ141 alloy. The results in Fig. 3(c) show that the Type C serrations occur at lower applied ε_ (curves 1 and 2) than do Type B serrations (curves 3–5), and that the Type C serrations exhibit stronger serrated stress drop/fluctuation than the Type B serrations. Also in Fig. 3(c), the type of serration changes from Type B to Type C at a transition point of E6% strain in curve 3, but at E12% strain in curve 4. Note that the ε_ for curve 4 (6.67
10  3 s  1) is twice that for curve 3 (3.33
10  3 s  1). The experimental results indicate that the transition time from Type B to Type C appearing in curves 3 and 4 is almost the same. However, further investigation is needed to clarify the observation shown in Fig. 3(c). Fig. 3(d) shows the relationship between flow stress (stress s at 0.2 strain) and ε_ from the results in Fig. 3(b). From Fig. 3(d), it is clear that the SRS, defined as ∂s0.2/∂ ln ε_ , is negative. The relationship between the negative SRS and the occurrence of the PLC effect is discussed in Section 3.4.

140

5

4

LZ141 HR ε =3.33x10-3 s-1

120

3

100

Stress σ (MPa)

36

120oC 70oC 25oC 0 oC -25oC -50oC

1 2 3 4 5 6

80 60

2 40

6

1

20 0 0

10

20

30

40

50

60

70

80

90

100

Strain ε (%)

3.3. The effect of temperature on the tensile stress–strain curve

3 4 5

120

100

Stress σ (MPa)

Fig. 4(a) shows the tensile stress–strain curves for HR LZ141 alloy tested at temperatures ranging from 120 1C to 50 1C with ε_ ¼3.33
10  3 s  1. Fig. 4(b) is the enlargement of the curves in Fig. 4(a) in the strain range 0–0.06. In Fig. 4(b), no serrated flow can be found in stress–strain curves 1, 2 or 3 at temperatures Z25 1C. However, for temperatures r0 1C, stress–strain curves 4–6 exhibit the PLC effect, with the serrated flow identified as Type A serrations, in which the serrations rise above the general level of the stress–strain curve and are periodic in nature [21]. Fig. 4(b) shows that the serrations in curves 4 and 5 at 0 1C and 25 1C, respectively, occur in a slowly serrated succession, but the speed of the serrated succession increases significantly in curve 6 at 50 1C. Fig. 4(b) also shows that, in curve 6 at  50 1C, the serrations start at the beginning of the stress–strain curve and finish at the ultimate tensile strength (at the beginning of the necking). Fig. 5(a) shows the tensile stress–strain curves of SS LZ141 alloy at temperatures ranging from 75 1C to  50 1C with ε_ ¼3.33
10  4 s  1. This strain rate was chosen because at ε_ ¼3.33
10  4 s  1, there is a clearly visible PLC effect in the stress–strain curve at room temperature, as shown in Fig. 3(c). Fig. 5(b) plots the enlargement of the curves in Fig. 5(a) in the strain range 0–0.05. In Fig. 5(b), no serrated flow can be observed in stress–strain curves 1 and 2 at 75 1C and 50 1C, respectively, but the PLC effect occurs in curves 3–6 at temperatures r25 1C. Fig. 5(b) also shows that for the occurrence of the PLC effect in the stress–strain curves, the serration density and its strain range increase as the testing temperature decreases. Type C serrations can be identified in curves 3 and 4, and Type B serrations appear in curves 5 and 6. In Fig. 5(a), there is a large decrement of the yielding strength in between 25 1C/0 1C tested curves and  25 1C/  50 1C tested curves. This feature is suggested to have connection with the change of the serration type shown in Fig. 5(b). But in curve 6, there is a transition point at E6% strain, where the serration type changes from Type B to Type C as the testing strain increases. Comparing Fig. 5(b) with Fig. 3(c), for the same Type B serrations, one can find that the serration density is much higher in Fig. 5(b). In Fig. 5(a) and (b), in curves 5 and 6, both serrations start to appear at the beginning of the stress–strain curve. Brindley et al. reported that, for Cu–Zn crystals, the stress–strain curve tested at 400 1C for Cu–5 at% Zn crystal and those tested at 200 1C and 400 1C for Cu–10 at% Zn crystal exhibited a fine serration throughout the test [24]. However, alloys that exhibit a serrated flow from the beginning of the tensile stress–strain curve are usually those tested at high temperatures, such as 200 1C and 400 1C, as

2

80

60

1 LZ141 HR -3 -1 ε=3.33x10 s

40

6

20

120oC 70oC 25oC 0 oC -25oC -50oC

1 2 3 4 5 6

0 0

1

2

3

4

5

6

Strain ε (%) Fig. 4. (a) Tensile stress–strain curves of hot-rolled (HR) LZ141 alloy tested at ε_ ¼ 3.33
10  3 s  1 at the temperatures of 120 1C to 50 1C. (b) Enlarged curves of (a) in the strain range 0–0.06.

reported in Ref. [24], but not those tested at low temperatures, such as  25 1C or 50 1C, as shown in Fig. 5(b). From curves 3–6 in Fig. 5(b), the regressive line of ln εc vs. 1/T is plotted and shown in Fig. 5(c), in which εc is the onset of the serration yielding in the stress–strain curve. In Fig. 5(c), the slope of the regressive line is negative. This characteristic is uncommon because for most alloys with a serrated flow in the stress–strain curve, the slope of the regressive line for the plot of ln εc vs. 1/T is positive, and this slope can be used to calculate the effective activation energy for the exchange of vacancies and solute atoms [9,10,21,25]. Brindley and Worthington found that for Al–3% Mg alloy tensile tested at ε_ ¼1.33
10  4 s  1 in the temperature range 70 1C to 150 1C, the plot of ln εc vs. 1/T has a negative slope from 25 1C to 150 1C, but this plot has a positive slope from 25 1C to  70 1C [26]. They discussed the activation energy associated with the serrated yielding for the tensile curves at temperatures from 25 1C to  70 1C, but the reason for the negative slope in the temperature from 25 1C to 150 1C was not discussed. Carefully examining the stress–strain curves shown in Fig. 5(a), one can see that the yielding strength of the SS alloy decreases dramatically as the testing temperature decreases. This characteristic is proposed to be closely related to the occurrence of the negative slope exhibited in Fig. 5(c). Fig. 5(c) also shows that the εc values for curves 3 and 4 at 25 1C and 0 1C, respectively, are significantly higher than those for curves 5 and 6 at  25 1C and  50 1C, respectively. In addition, the serration changes from Type C to

S.K. Wu et al. / Materials Science & Engineering A 605 (2014) 33–38

120

120

1 2 3

LZ141 SS ε=3.33x10-4 s-1

100

75 C 4 50 C 5 25 C 6

0C -25 C -50 C

37

1 2 3

LZ141 SS -4 -1 ε=3.33x10 s

75 C 4 50 C 5 25 C 6

0C -25 C -50 C

3

100

Stress σ (MPa)

Stress σ (MPa)

2 80

3

60

2 40

6

5

1

4

80

1 4

60

40

6 5

20

20

0

0 0

20

40

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0

1

2

3

4

5

6

Strain ε (%)

Strain ε (%)

70

-3

LZ141 SS ε=3.33x10-4 s-1

4 -4

3

LZ141 SS ε=3.33x10-4 s-1

60

5 6

-25 C -50 C

Stress σ (MPa)

50

ln εc

-5

-6

40 30

6

5

20

-7

5

10

6

-8

0

3.2

3.4

3.6

3.8

4.0

4.2

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4.6

-1

1000/T(K )

0

5

10

15

20

25

30

Strain ε (%)

Fig. 5. (a) Tensile stress–strain curves of solid-solution treated (SS) LZ141 alloy tested at ε_ ¼3.33
10  4 s  1 at temperatures of 75 1C to 50 1C. (b) Enlarged curves of (a) in the strain range 0–0.05. (c) The plot of the ln εc vs. 1/T for the curves 3–6 is shown in (b). (d) Magnified curves 5 and 6 of (b), in which the arrows indicate the locations of the large stress variations appearing on the curves.

Type B as the testing temperature decreases. These results are consistent with the characteristics of the PLC effect reported in Refs. [21,26]; however, the alloys studied in Refs. [21,26] exhibit a positive slope in the plot of ln εc vs. 1/T, but the SS LZ141 alloy shown in Fig. 5(c) has a negative slope. Fig. 5(d) is an enlargement of Fig. 5(b) in the strain range 0–0.30 for curves 5 and 6 at 25 1C and  50 1C, respectively. In Fig. 5(d), serrations with a stress variation of about 3 MPa are indicated by arrows. Fig. 6 is the OM image of the  50 1C tensile specimen of the SS alloy taken near the fractured surface. In Fig. 6, only the slip bands can be found in the elongated grains and no twins can be observed. This indicates that the DSA mechanism, not the twinning mechanism, causes the occurrence of the PLC effect shown in Fig. 5(d). The twinning mechanism of the PLC effect has been reported for Mg–5Li–3Al–1.5Zn–2RE alloy with an HCP structure under ε_ Z1
10  3 s  1, in which severe serrations with a stress variation of about 20 MPa are found [27].

100μm Fig. 6. Optical micrograph taken near the fractured surface of the tensile specimen of SS LZ141 alloy tested at  50 1C. The fractured surface is located in the upper-left corner.

3.4. The characteristics of the serrated flow in HR and SS LZ141 alloys From Figs. 1 and 2, it is reasonable to recognize that the dislocation density is higher and the number of solute atoms is lower in the HR alloy than in the SS alloy. In Fig. 3(a), for the HR alloy, the occurrence of the PLC effect is not observed in tensile curves for all ε_ . However, for the SS alloy in Fig. 3(b), the PLC effect is observed in tensile curves 1–5. The DSA mechanism can explain the occurrence of the PLC effect observed in Fig. 3

(a) and (b), as mentioned in Section 3.2. Moreover, Fig. 3 (d) indicates that the serrated flow in the SS alloy has a negative SRS. In a study of LA41 magnesium alloy with an HCP structure, Wang et al. reported that, a serrated flow with a negative SRS can be explained by the DSA mechanism [10]. Therefore, the result in Fig. 3(d) also supports the proposition that the DSA mechanism causes the occurrence of the PLC effect in the SS alloy.

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S.K. Wu et al. / Materials Science & Engineering A 605 (2014) 33–38

Fig. 3 shows that the effect of the strain rate ε_ on the occurrence of the PLC effect is significant in the SS alloy. Figs. 4 and 5 indicate that the effect of the temperature T on the occurrence of the serrated flow is remarkable in both HR and SS alloys. In Figs. 4 and 5, for both HR and SS alloys at temperatures r25 1C, the serration density increases as the testing temperature decreases. Comparing Fig. 3 (c) with Fig. 5(b), for the occurrence of the PLC effect in the SS alloy, one can see that a decrease in T is equivalent to an increase in ε_ . For example, for the SS alloy, the change of the serrated flow from curve 3 at 25 1C to curve 4 at 0 1C (Fig. 5(b)) is quite similar to the change from curve 1 at ε_ ¼ 3.33
10  4 s  1 to curve 2 at ε_ ¼ 6.67
10  4 s  1 (Fig. 3(c)). This phenomenon has also been found in Cu–1.09 at% In alloy tested at temperatures of 20–400 1C in the ε_ range of around 10  4 s  1 [21]. It is also noted that, for temperatures r 25 1C, the serrated flow in the HR alloy (Fig. 4) is identified as Type A serrations, but that the SS alloy (Fig. 5) has Type B and Type C serrations, as previously discussed in Section 3.3. The change of the serrations type shown in Figs. 4 and 5 can be explained by the viewpoint of Ref. [23]. That is, from Figs. 1 and 2, it appears that the HR alloy has an intrinsically higher dislocation density and fewer solute atoms than the SS alloy. These characteristics cause a weak DSA effect in the HR alloy but induce an enhanced DSA effect in the SS alloy; therefore, different types of serrations can be seen in Figs. 4 and 5. For the SS alloy, Fig. 5(d) shows Type B serrations in curve 5, and they also occur in curve 6 at lower strain (ɛ o E 6%). Fig. 5(d) also shows that in curve 6, the serration type changes from Type B to Type C when the strain is higher than E 6%, as discussed in Section 3.3. This characteristic comes from the fact that the mobility of the dislocations decreases as the temperature decreases, and the multiplication of the dislocations is enhanced at higher strain; thus, an enhanced DSA effect with Type C serrations appears. 4. Conclusions The PLC effect has been observed in tensile HR and SS β-phase LZ141 magnesium alloys for strain rates (ε_ ) from 3.33
10  4 s  1 to 3.33
10  2 s  1 and at temperatures (T) from 25 1C to 50 1C. Both ε_ and T affect the occurrence of the PLC effect. The HR alloy has an intrinsically higher dislocation density and fewer solute atoms than the SS alloy. The dynamic strain aging (DSA) mechanism explains why the PLC effect does not occur in the HR alloy at room temperature but does occur in the SS alloy, which exhibits Type B serrations at ε_ ¼(3.33–6.67)
10  3 s  1 in the lower strain range and Type C serrations at ε_ ¼(3.33–6.67)
10  3 s  1 in the higher strain range and at ε_ ¼(3.33–6.67)
10  4 s  1. At the same time, for the SS alloy at room temperature, the strain rate sensitivity (SRS) is negative. The negative SRS supports the proposition that the DSA mechanism is responsible for the occurrence of the PLC effect. In the study of the T effect on the appearance of the serrated flow, for the HR alloy at T r0 1C, Type A serrations were apparent, and for the SS alloy, Type C serrations occurred in the curves at 25 1C and 0 1C. Type B serrations

occurred in curves at  25 1C and  50 1C, while in the curve at  50 1C, the serration type changed from Type B to Type C at high strain. These results can also be explained by the DSA mechanism. In the SS alloy, the effect of a decrease in T from 25 1C to 0 1C at 3.33
10  4 s  1 is equivalent to an increase in ε_ from 3.33
10  4 s  1 to 6.67
10  4 s  1 at 25 1C. In addition, large stress variations appeared in the serrated flows of the curves at  25 1C and  50 1C for the SS alloy, but no twinning was observed near the fractured surface. Experimental results indicate that for LZ141 alloy at low temperatures, the PLC effect starts at the beginning of the stress–strain curves and that the slope of the regressive line for the plot of ln εc vs. 1/T is negative. Both phenomena are uncommon and require further investigation.

Acknowledgment The authors gratefully acknowledge the financial support for this study from the National Science Council Taiwan under Grants NSC1002221-E002-102-MY3 and NSC102-2623-E259-001-D. C. Chien also acknowledges the support for this research from the National Taiwan University (NTU) under Grant no. NTU102R891803 (The Excellence Research Program, NTU). References [1] ASM. Metals Handbook, Vol. 3, Alloys Phase Diagrams, ASM International, Ohio, USA, 1992, pp. 2. 276. [2] W.R.D. Jones, J. Inst. Met. 84 (1955–56) 364–378. [3] S.K. Wu, Y.H. Li, K.T. Chien, C. Chien, C.S. Yang, J. Alloys Compd. 563 (2013) 234–241. [4] S.L. Couling, Acta Metall. 7 (1959) 133–134. [5] K.G. Kent, A. Kelly, J. Inst. Met. 93 (1964–65) 536–542. [6] M. Chaturvedi, D.J. Lloyd, K. Tangri, Met. Sci. 6 (1972) 16–19. [7] M.C. Chaturvedi, D.J. Lloyd, Philos. Mag. 30 (1974) 1199–1207. [8] C. Corby, C.H. Cáceres, P. Lukáč, Mater. Sci. Eng. A 387-389 (2004) 22–24. [9] S.M. Zhu, J.F. Nie, Scr. Mater. 50 (2004) 51–55. [10] C. Wang, Y.B. Xu, E.H. Han, Mater. Lett. 60 (2006) 2941–2944. [11] A.H. Cottrell, Philos. Mag. 44 (1953) 829–832. [12] J.B. Clark, L. Sturkey, J. Inst. Met. 86 (1957–1958) 272–276. [13] D.W. Levinson, D. McPherson, Trans. ASM 48 (1956) 689–705. [14] H. Takuda, H. Matsusaka, S. Kikuchi, K. Kubota, J. Mater. Sci. 37 (2002) 51–57. [15] G.S. Song, M. Staiger, M. Kral, Mater. Sci. Eng. A 371 (2004) 371–376. [16] T. Liu, S.D. Wu, S.X. Li, P.J. Li, Mater. Sci. Eng. A 460–461 (2007) 499–503. [17] S.K. Wu, J.Y. Wang, K.C. Lin, H.Y. Bor, Mater. Sci. Eng. A 522 (2012) 76–80. [18] R.W.K. Honeycombe, The Plastic Deformation of Metals, E. Arnold, London (1984) 150–151. [19] ASTM E112-88, Standard Test Methods for Determining Grain Size, ASTM International. [20] ASTM E8/E8M-13a, Standard Test Methods for Tension Testing of Metallic Materials, ASTM International. [21] B.J. Brindley, P.J. Worthington, Metall. Rev. 15 (1970) 101–114. [22] K. Chihab, Y. Estrin, L.P. Kubin, J. Vergnol, Scr. Metall. Mater. 21 (1987) 203–208. [23] H. Jiang, Q. Zhang, X. Chen, Z. Chen, Z. Jiang, X. Wu, J. Fan, Acta Mater. 55 (2007) 2219–2228. [24] B.J. Brindley, D.J.H. Corderoy, R.K. Honeycombe, Acta Metall. 10 (1962) 1042–1050. [25] P.G. McCormick, Acta Metall. 19 (1971) 463–471. [26] B.J. Brindley, P.J. Worthington, Acta Metall. 17 (1969) 1357–1361. [27] T.Q. Li, Y.B. Liu, Z.Y. Cao, R.Z. Wu, M.L. Zhang, L.R. Cheng, D.M. Jiang, J. Alloys Compd. 509 (2011) 7607–7610.