Creep and fracture behavior of as-cast Mg–11Y–5Gd–2Zn–0.5Zr(wt%)

D. D. Yin • Q. D. Wang • C. J. Boehlert •

W. J. Ding

Received: 7 February 2012 / Accepted: 30 April 2012 / Published online: 1 June 2012

� Springer Science+Business Media, LLC 2012

Abstract The tensile-creep and creep–fracture behavior

of as-cast Mg–11Y–5Gd–2Zn–0.5Zr (wt%) (WGZ1152)

was investigated at temperatures between 523 and 598 K

(0.58–0.66Tm) and stresses between 30 and 140 MPa. The

creep stress exponent was close to five, suggesting that

dislocation creep was the dominant creep mechanism. The

activation energy for creep (233 ± 18 kJ/mol) was higher

than that for self-diffusion in magnesium, and was believed

to be associated with cross-slip, which was the dominant

thermally-aided creep mechanism. This was consistent

with the surface observations, which suggested non-basal

slip and cross-slip were active at 573 K. The minimum

creep rate and fracture time values fit the original and

modified Monkman–Grant models. In situ creep experi-

ments highlighted the intergranular cracking evolution. The

creep properties and behavior were compared with those

for other high-temperature creep-resistant Mg alloys such

as WE54-T6 and HZ32-T5.

Introduction

Due to their low density, high specific strength, specific

stiffness, damping capacity, good castability, machinabil-

ity, and recyclability, magnesium (Mg) alloys are promis-

ing light-weight structural material for applications in

which weight reduction is critical, including aerospace,

aircraft, and automotive components [1–5]. However, their

limited strength and creep resistance at temperatures above

398 K have limited their use [2–4, 6].

Al is the most widely used alloying element in Mg

alloys due to its strengthening effect. Although Mg–Al–Mn

(AM) and Mg–Al–Zn (AZ) series alloys are strong at room

temperature, they exhibit relatively poor creep resistance at

temperatures above 393–423 K compared with rare-earth

(RE) containing Mg alloys due to the instability of the

b-Mg17Al12 phase [4]. Small additions of Si, RE elements,

and Ca to AZ and AM series alloys can increase the creep

resistance, as has been shown for the following alloys:

AS21, AS41, AE42, AE41, AJ52, AX53, and MRI153

[3–5, 7, 8]. Ag and RE elements are effective creep

strengtheners at temperatures up to 523 K, and this finding

has resulted in the development of Mg–Ag–RE–Zr alloys

including QE22 [5, 6]. Even greater creep resistance can be

achieved through the addition of Th or RE (Y and Gd) to

Mg [3, 5, 6, 9]. In this group, HZ32 [Mg–3.3Th–2.1Zn–

0.7Zr (wt%)] is the Mg alloy which has been used at the

highest service temperatures (618–643 K) [3, 6]. However,

HZ32 is being phased out due to its radioactivity [3].

WE54 [Mg–5.2Y–3.6RE–0.5Zr (wt%)] is the most widely

used commercial structural Mg alloys for elevated-tem-

perature applications up to 523 K [1, 6, 10]. Recently, a

Mg–Y based alloy, Mg–11Y–5Gd–2Zn–0.5Zr (wt%)

(WGZ1152), has been developed, and this alloy has

exhibited potential for elevated-temperature applications at

D. D. Yin � Q. D. Wang (&) � W. J. Ding

National Engineering Research Center of Light Alloy Net

Forming, School of Materials Science and Engineering,

Shanghai Jiao Tong University, 800 Dongchuan Road,

Shanghai 200240, People’s Republic of China

e-mail: wangqudong@sjtu.edu.cn

D. D. Yin

e-mail: ahnydd@gmail.com

Q. D. Wang � W. J. Ding

The State Key Laboratory of Metal Matrix Composites,

School of Materials Science and Engineering, Shanghai Jiao

Tong University, 800 Dongchuan Road, Shanghai 200240,

People’s Republic of China

C. J. Boehlert

Department of Chemical Engineering and Materials Science,

Michigan State University, East Lansing, MI 48824, USA

123

J Mater Sci (2012) 47:6263–6275

DOI 10.1007/s10853-012-6546-4

temperatures above 523 K [11–14]. In order to further

improve and develop high-performance creep-resistant Mg

alloys, it is necessary to systematically investigate their

mechanical behavior and deformation mechanisms. The

aim of the present work was to investigate the tensile-creep

and creep–fracture behavior of as-cast WGZ1152 at

523–598 K (0.58–0.66Tm, where Tm is the absolute melting

temperature of this alloy [12]). The relationship among

microstructural evolution, creep deformation, and creep–

fracture are discussed.

Experimental procedures

The measured composition of the studied alloy, measured

using an inductively coupled plasma analyzer, was Mg–

11.3Y–4.7Gd–2.0Zn–0.46Zr (wt%). It was prepared from

high-purity Mg, Zn ingots (C99.9 %), and Mg–25Y (wt%),

Mg–25Gd (wt%), and Mg–30Zr (wt%) master alloys by

electronic resistance melting in a mild steel crucible at

approximately 1,023 K under a mixed atmosphere of CO2

and SF6. The ratio of CO2 gas to SF6 gas was 100:1. The

alloy was cast into a rectangular steel mold with dimen-

sions of 210 mm in length, 130 mm in width, and 60 mm

in thickness, which was pre-heated to 473 K. The

mechanical properties of the WGZ1152 alloy were com-

pared with that for a WE54 [Mg–5.2Y–3.6RE–0.5Zr

(wt%)] ingot which was purchased from Magnesium

Elektron UK and subjected to a solution treatment at 798 K

for 8 h followed by water quenching and a subsequent

aging treatment of 16 h at 523 K [1].

Conventional tensile-creep tests were performed at tem-

peratures between 523 and 598 K and under constant load in

air. Rectangular specimens, with dimensions of 25 mm in

gage length, 6 mm in width, and 2 mm in thickness, were cut

from the ingot by an electric discharge machine (EDM). In

order to achieve thermal stability, the specimens were held at

the test temperature for 1 h before loading. The temperature

was maintained constant within ±1 K of the target temper-

ature throughout the tests. The strain was measured using a

linear variable differential transformer attached to a high-

temperature extensometer, which was attached to the gage

section of the specimens. The temperature-stress–strain-time

data were collected by a computer-controlled data acquisi-

tion system. Most of the specimens were taken to failure.

Some of the tests were interrupted after 1,000 h or longer.

An in situ tensile-creep experiment was conducted on

the as-cast material at T = 573 K and r = 50 MPa. A flat

dogbone-shaped sample, with dimensions of 10 mm in

gage length, 3 mm in width, and 2.5 mm in thickness, was

cut using EDM. The specimen was glued to a metallic

platen and polished to a 1 lm finish using an automatic

polishing machine and ethanol as a polishing lubricant. The

experiment was performed using a screw-driven tensile

stage (built by Ernest F. Fullam, Incorporated, Clifton

Park, NY) placed inside a Zeiss (Jena, Germany) EVO

MA15 scanning electron microscope (SEM). Temperature

was controlled using a constant-voltage power supply to a

6 mm diameter tungsten-based heater located just below

the gage section of the sample. An open-bath, closed-loop

chiller was used to circulate distilled water at RT through

copper tubes to prevent the tensile stage from overheating.

A fine-gage type-K thermocouple was spot-welded to the

gage section of the sample. After the sample’s gauge-sec-

tion temperature reached the desired creep temperature, a

30 min period was given to stabilize the thermal stress

prior to applying load. The load, which was measured

using a 4,448 N load cell, was applied at 5 N/s until

reaching the desired creep stress. The test was considered

constant load where the stress fluctuation varied ±3 MPa.

Secondary electron (SE) SEM images were taken before

loading and at periodic displacements throughout the

experiment without interrupting the experiment. The

pressure in the SEM chamber never exceeded 10-6 torr,

and therefore, oxidation did not detrimentally affect the

SEM imaging. Further details of this apparatus and testing

technique can be found elsewhere [15, 16].

Microstructural observations, both before and after

deformation, were made using optical microscopy (LEICA

MEF4M), and scanning electron microscopy (FEI SIRION

200). The creep specimens, for which ex-situ observations

obtained, were mechanically polished, and then etched in a

4 vol% nital solution.

Results

Initial microstructure

The microstructure of the WGZ1152 has been described

previously [11, 12]. Figure 1a illustrates a representative

optical photomicrograph of the as-cast alloy. The gray

Mg24(GdYZn)5 phase and the strip-shaped black X phases

were distributed along the grain boundaries. The fine-lamellar

phases were uniformly distributed from grain boundaries to

the interior of the a-Mg grains and exhibited one single ori-

entation in a given grain as shown in Fig. 1b. Based on pre-

vious studies [11, 12, 17], both the X phase along grain

boundaries and the fine-lamellar phases within grain interior

exhibited an 18R-type long period stacking ordered structure.

The average a-Mg grain size was 116 ± 15 lm [12].

Creep properties

Figure 2 compares the minimum creep rates of the studied

alloy with that of a WE54-T6 alloy and a HZ32-T5 alloy.

6264 J Mater Sci (2012) 47:6263–6275

123

The minimum creep rate of the as-cast WGZ1152 was

superior to the WE54-T6 at T = 523–573 K (0.58–0.64Tm)

and r = 50–80 MPa, and was comparable to the HZ32-T5

at T = 573 K and r = 30 MPa. At T = 523 K and

r = 80 MPa, the minimum creep rate of the as-cast

WGZ1152 was almost two orders of magnitude lower than

that for WE54-T6. At T = 573 K and r = 50 MPa, the

minimum creep rate of the as-cast WGZ1152 was almost

one order of magnitude lower than that for WE54-T6, and

the creep–fracture life of the as-cast WGZ1152 (699 h)

was 12 times higher than that for WE54-T6 (*53 h). The

minimum creep rate of the as-cast WGZ1152 was

comparable to that for HZ32-T5 at T = 573 K and

r = 30 MPa [10].

Figure 3 illustrates representative creep curves of as-

cast WGZ1152 at T = 573 K and r = 30–80 MPa. The

creep resistance decreased with increasing stress. The

corresponding creep strain rate versus creep strain curves

are illustrated in Fig. 4. After the primary creep stage, the

creep strain rate reached a minimum, and then increased

quite slowly with increasing strain. In the tertiary creep

regime, the creep strain rate increased with increasing

strain until fracture occurred.

Figure 5 shows the log–log plot of the stress dependence

on the minimum creep rate at T = 523–598 K. The creep

stress exponent n was determined from the slope. For 523,

573, and 598 K the n values were 5.0, 4.6 and 5.1,

respectively. The Arrhenius plot of the logarithm of the

minimum creep rate versus the reciprocal of temperature at

stresses between 50 and 140 MPa is illustrated in Fig. 6.

The activation energy for creep was determined from the

slope. The activation energy for creep was independent of

stress within the experimental deviation (Fig. 7). The

average activation energy for creep Qavg was 233 ±

18 kJ/mol for all the stress levels examined.

Fig. 1 a Optical photomicrograph and b secondary electron SEM photomicrograph of the as-cast WGZ1152 showing its representative

microstructure

Fig. 2 Minimum creep rates of the as-cast WGZ1152 in comparison

with WE54-T6 and HZ32-T5

0 200 400 600 800 1000 12000

2

4

6

8

10

12

did not fractureCre

ep S

trai

n (

%)

Time (h)

573K/30MPa 573K/50MPa 573K/80MPa

Fig. 3 Representative creep curves of the as-cast WGZ1152

J Mater Sci (2012) 47:6263–6275 6265

123

Creep fracture

The creep damage tolerance parameter k measures the

tolerance of a material to strain concentrations, and is given

by [18].

k ¼ er

_emintr: ð1Þ

The plot of er versus _emintr is illustrated in Fig. 8. The creep

damage tolerance parameter k was determined from the

slope. The k values ranged from 1.2 to 2.3 for most

test conditions. Only two test conditions, which were

T = 573 K/r = 80 MPa and T = 573 K/r = 100 MPa,

exhibited a higher k value (k & 2.7). Theoretical calcula-

tion has predicted that when k ranges from 1 to 2.5, grain

boundary cavities and cracks attribute significantly to ter-

tiary creep and final fracture [19]. In the present study, the

k range of 1.2 to 2.3 suggested that grain boundary cavities

and cracks may play an important role in creep fracture.

Figure 9 illustrates the creep–fracture properties at

T = 523–598 K. There was no definite trend in the fracture

strain with respect to stress at all the temperatures evalu-

ated (Fig. 9a), however the fracture time decreased with

increasing stress (Fig. 9b).

1E-8

1E-7

1E-6

1E-5

Cre

ep S

trai

n R

ate

(s-1) 573K/50MPa

573K/80MPa

t/tr

1E-9

1E-8

1E-7

1E-6

1E-5

Cre

ep S

trai

n R

ate

(s-1)

Creep Strain (%)

573K/50MPa 573K/80MPa

(a) (b)

1E-9

1E-8

1E-7

1E-6

1E-5

did not fracture

Cre

ep S

trai

n R

ate

(s-1)

Time (h)

573K/30MPa

0.0 0.2 0.4 0.6 0.8 1.0 0 2 4 6 8 10 12

0 200 400 600 800 1000 0.0 0.2 0.4 0.61E-9

1E-8

1E-7

1E-6

1E-5

did not fracture

Cre

ep S

trai

n R

ate

(s-1)

Creep Strain (%)

573K/30MPa(d)(c)

Fig. 4 Representative creep

rate curves of the as-cast

WGZ1152: a and c creep strain

rate versus time; b and d creep

strain rate versus creep strain.

Note that in Fig. 3a the creep

time was normalized by the

fracture time, tr, for the

specimens which failed

160140120100806040201E-10

1E-9

1E-8

1E-7

1E-6

1E-5

1E-4

n=5.0

n=4.6

n=5.1

598K 573K 523K

Min

imu

m C

reep

Rat

e (s

-1)

Stress (MPa)

Fig. 5 Log–log plot of stress dependence on the minimum creep rate

at T = 523–598 K

0.00165 0.00170 0.00175 0.00180 0.00185 0.00190 0.00195

1E-9

1E-8

1E-7

1E-6

1E-5

523K

573K

598K

50MPa 80MPa 100MPa 120MPa 140MPa

Min

imu

m C

reep

Rat

e (s

-1)

1/T (K-1)

Fig. 6 The Arrhenius plot of logarithm of minimum creep rate versus

the reciprocal of temperature at r = 50–140 MPa

6266 J Mater Sci (2012) 47:6263–6275

123

Under intergranular creep–fracture conditions, it is often

observed that the minimum creep rate and the fracture time

for many metals and alloys can be related by the Monk-

man–Grant (MGR) equation [20, 21]

_emintmr ¼ C ð2Þ

where m, and C are materials constants. The data for many

materials is better fitted by the modified Monkman–Grant

relationship (MMGR) [22]

tr _em0

min

er

¼ C0 ð3Þ

where m0

is close to unity and C0

is independent of tem-

perature and stress. The linear relationship of minimum

creep rate and time-to-fracture in log–log coordinates

shows the MGR relationship according to Eq. 2 (Fig. 10).

The m and C values were 1.04 and 0.02, respectively. The

linear relationship of mean creep rate er=tr versus minimum

creep rate in log–log coordinates shows MMGR according

to Eq. 3 (Fig. 11). The m0

and C0

values were 0.92 and

1.55, respectively. The m, C, m0, and C

0values were in the

common range observed for most of metals and alloys [21,

23–25].

Microstructural evolution during creep exposure

Figure 12 illustrates precipitates formed during the creep

exposure at 573 K. After 8 h of creep exposure at

T = 573 K and r = 120 MPa (Fig. 12a), no obvious

changes were detected compared with the initial micro-

structure. After 35 h of creep exposure at T = 573 K and

r = 100 MPa (Fig. 12b), dense spheroid precipitates

(average diameter 0.5–1 lm) preferentially formed near

grain boundaries. After 252 h of creep exposure at

T = 573 K and r = 50 MPa, plate-shaped precipitates

(average length 4–10 lm) were dominant. They were also

preferentially located near grain boundaries, see Fig. 12c,

d. Besides the dominated plate-shaped large precipitates, a

few spheroid small precipitates were present. Taking into

account the four-stage precipitation sequence: a-Mg solid–

solution ? b00 (Do19) ? b0 (bco) ? b1 (fcc) ? b (fcc)

in the Mg–RE alloys [26, 27], it is likely that the spheroid

precipitates were the metastable b0 phase and the

plate-shaped precipitates were the equilibrium b phase.

40 60 80 100 120 1400

50

100

150

200

250

300

Qavg=233±18kJ/mol

Act

ivat

ion

En

erg

y Q

(kJ

/mo

l)

Stress (MPa)

Fig. 7 Activation energy Q for creep versus applied stress

0.00 0.01 0.02 0.03 0.04 0.05 0.060.00

0.04

0.08

0.12

0.16

0.20

3λ =

2.5λ =

1λ =

598K 573K 523K

Fra

ctu

re S

trai

n

min rt⋅

Fig. 8 Relationship between fracture strain er and _emin � tr, showing

the creep damage tolerance parameter k at T = 523–598 K

150100500

5

10

15 598K 573K 523K

Fra

ctu

re S

trai

n (

%)

Stress (MPa)15010050

1

10

100

1000 598K 573K 523K

Fra

ctu

re T

ime

(h)

Stress (MPa)

(a) (b)Fig. 9 Creep–fracture

properties at T = 523–598 K.

Stress dependence of a fracture

strain and b fracture time

J Mater Sci (2012) 47:6263–6275 6267

123

Therefore, it is suggested that the dominant precipitates

changed from b0 phase to b phase with increasing creep

time.

Figures 13 and 14 (low and high magnification) illus-

trate the deformation evolution during the in situ creep test

at T = 573 K and r = 50 MPa. No obvious cracking was

observed in the primary creep regime (Fig. 14b) and early

stages of the secondary creep regime (Fig. 13b). Cracking

became obvious within the later stages of the secondary

creep regime (Fig. 14c). Cracking initiated preferentially at

grain boundaries and near second phases located along the

grain boundaries. The crack propagation tended to follow

the grain boundaries. With increased creep time, the

intergranular cracks grew in length and width. The offsets

of fiducial scratches on the specimen surface were

observed at grain boundaries (Fig. 14, indicated with

arrows), and with increased creep time, the offset distances

became larger. This indicated that grain boundary sliding

may play an important role in the creep deformation and

fracture. Some short wavy slip traces were observed on the

surface of the sample in situ creep tested at T = 573 K and

r = 50 MPa (Fig. 15). Although the slip systems were not

identified for these slip traces, the wavy slip trace suggests

that non-basal slip and cross-slip were active [28, 29].

Discussion

Creep deformation mechanisms

At elevated temperatures, the minimum creep rate (_emin) of

many metals and alloys is generally described by the fol-

lowing equation [4, 28–31]

_emin ¼ Arn expð�Q=RTÞ ð4Þ

where A is a material constant, r is the stress, n is the stress

exponent, Q is the activation energy for creep, R is the gas

constant and T is the absolute temperature. Based on Eq. 4,

the n and Q can be defined by

n ¼ ½o ln _emin

o ln r�T ð5Þ

Q ¼ ½ o ln _emin

oð�1=RTÞ�r ð6Þ

the n and Q are sometimes useful for determining the

dominant creep deformation mechanisms empirically.

As seen in Fig. 5, the n was close to 5 for all the test

temperatures examined implying that dislocation creep

was the creep mechanism [30, 32]. For dislocation creep

[4, 30, 32], it is generally accepted that the flow stress is

thermal and athermal components. The athermal compo-

nent is a long-range component of flow stress. The ther-

mal component is local and thermal activated, and it can

help the dislocations overcome the obstacles effectively

during creep. Viscous glide, dislocation climb, cross-slip,

and dislocation jogs are important thermally-aided dislo-

cation creep mechanisms [4]. It should be noted that

different mechanisms may operate simultaneously.

Table 1 summaries the creep stress exponent (n), activa-

tion energy for creep (Q), and proposed creep deforma-

tion mechanisms for pure Mg [28, 29, 33, 34] and some

Mg alloys, including Mg–Al based alloys [28, 29, 35–39],

Mg–Zn based alloys [40, 41], Mg–Y based alloys [10, 24,

42–47], Mg–Gd based alloys [9, 10, 44], and Mg–Sc

based alloys [10]. As shown in table 1, an n value of *5

is commonly observed for Mg and Mg alloys and the

creep mechanism is believed to be dislocation creep. The

1000 10000 100000 1000000 1E71E-9

1E-8

1E-7

1E-6

1E-5

1E-4

598K 573K 523K

Min

imu

m C

reep

Rat

e (s

-1)

Fracture Time (s)

min

m

rt Cε•

=

2

1.04

0.02

. 0.95

m

C

Adj R

==

=

Fig. 10 Log–Log plot of minimum creep rate versus fracture time at

T = 523–598 K and stresses between 50 and 140 MPa showing

Monkman–Grant relationship (MGR)

1E-8 1E-7 1E-6 1E-5 1E-41E-8

1E-7

1E-6

1E-5

1E-4 598K 573K 523K

r/tr (

s-1)

Minimum Creep Rate (s-1)

'

'minm

r

r

tC

ε⋅ =

'

'

2

0.92

1.55

. 0.97

m

C

Adj R

==

=

Fig. 11 Log–Log plot of er=tr versus minimum creep rate at

T = 523–598 K and stresses between 50 and 140 MPa showing

modified Monkman–Grant relationship (MMGR)

6268 J Mater Sci (2012) 47:6263–6275

123

Q value creep can provide more information about the

specified thermally-aided dislocation creep mechanisms.

The measured Qavg value of the studied alloy was

233 ± 18 kJ/mol, and this value was significantly higher

than the activation energy for lattice self-diffusion of pure

Mg (135 kJ/mol) [1]. Similar Q values have been reported

by Vagarali et al. [28, 29], Milika et al. [33], Suzuki et al.

[42, 43], and Regev et al. [37, 38]. Vagarali et al. [28]

performed tensile creep tests for pure Mg and found that

the activation energy for creep was 190–220 ± 10 kJ/mol

for temperature between 750 and 820 K and stresses

between 3.7 and 6.3 MPa, and there was evidence of non-

basal slip. They compared the cross-slip mechanisms

developed by Friedel with their experimental data, and

concluded that the creep deformation was controlled by

the rate of cross-slip of dislocations from the basal to the

prismatic planes. They also studied the creep behavior of

a M–0.8Al solid–solution alloy, and found a Q value of

230 ± 15 kJ/mol at temperatures of 750–810 K and

stresses below 10 MPa [29]. The most important micro-

structural features were extensive non-basal slip traces

and some isolated, but well-defined sub-boundaries. As in

their previous study for pure Mg, the experimental data

were in agreement with a cross-slip mechanism. Thus,

they concluded the cross-slip of screw dislocations from

the basal to the prismatic planes was the rate-controlling

mechanism. Milika et al. [33] reported an activation

energy for creep of 160–240 kJ/mol for pure Mg at

temperatures of 400–800 K and stresses below 30 MPa.

They suggested the cross-slip of screw dislocations from

the basal to the pyramidal planes and pyramidal slip

limited by nucleation of kink motion were the possible

dislocation creep mechanisms. Suzuki et al. [42, 43]

studied the creep behavior of binary Mg–(0.7–3.9)Y

alloys at temperatures between 550 and 650 K and

stresses between 4 and 40 MPa. An Q value of

230–290 kJ/mol was reported, and cross-slip of screw

\a[-dislocations were observed at T = 650 K. They

believed that the relatively high activation energy was

associated with cross-slip. Regev et al. [37, 38] studied

the creep behavior of AZ91D at temperatures between

392 and 453 K and stresses between 40 and 115 MPa.

Their Q vaule was 94–220 kJ/mol, and non-basal dislo-

cation segments were identified using TEM. They pro-

posed that cross-slip might control the creep deformation.

Thus, it is reasonable to believe that the activation energy

Fig. 12 a–c Optical photomicrographs and d Secondary electron

SEM photomicrograph of specimens creep tested to different creep

times: a 8 h (T = 573 K and r = 120 MPa), b 35 h (T = 573 K and

r = 100 MPa), c and d 252 h (T = 573 K and r = 50 MPa). The

tensile axis was horizontal

J Mater Sci (2012) 47:6263–6275 6269

123

for creep measured in the current study was influenced by

cross-slip.

Creep–fracture mechanisms

The in situ creep experiment indicated that intergranular

fracture was the dominant creep–fracture mode. This

observation is consistent with both the range of the creep

damage tolerance parameter k and the original and MMGR.

Intergranular fracture involves creep cavity nucleation

followed by coalescence and growth of isolated cavities

and their linkage to form cracks [48, 49]. Eventually this

leads to final failure. Creep cavity nucleation and

growth requires sufficient stress concentration [48]. The

microhardness of the grain boundary second phases,

Mg24(GdYZn)5 and X phase, were 32 and 45 % larger,

Fig. 13 Secondary electron SEM photomicrographs (low magnifica-

tion) of a specimen taken during an in situ creep test at T = 573 K

and r = 50 MPa, showing intergranular cracking evolution. The

creep stages and approximate strain values are indicated in the

figures. The tensile axis was horizontal

6270 J Mater Sci (2012) 47:6263–6275

123

respectively, than that of Mg matrix, and their volume

fraction was more than 9 % [12]. These hard and numerous

grain boundary second phases may have assisted disloca-

tion pile-ups and generated sufficient stress concentrations

for cavity nucleation and growth. Another important factor

which may aid creep cavity nucleation and growth was

grain boundary sliding (Fig. 15). This can lead to sufficient

stress concentration at triple points and hard particles to

form grain boundaries cavities [48]. Also, according to the

model proposed by Chen [50], grain boundary sliding can

prompt cavity growth by enhancing effective mass trans-

port rates along the sliding grain boundary.

Conclusions

The detailed tensile-creep and creep–fracture behavior

of as-cast WGZ1152 was investigated at temperatures

Fig. 14 Secondary electron SEM photomicrographs (high magnifi-

cation) of a specimen taken during an in situ creep test at T = 573 K

and r = 50 MPa, showing deformation evolution. The creep stages

and approximate strain values are indicated in the figures. The tensile

axis was horizontal. The arrows are pointing to the offsets of fiducial

scratches on the specimen surface at grain boundaries

J Mater Sci (2012) 47:6263–6275 6271

123

between 523-598 K (0.58-0.66Tm) and stresses between

30 MPa to 140 MPa. The main conclusions of this work

are summarized below:

(1) The minimum creep rate for as-cast WGZ1152 was

more than one magnitude lower than that for a

WE54-T6 alloy at T = 523–573 K (0.58–0.64Tm)

Fig. 15 Secondary electron SEM photomicrograph of the surface of an in situ creep tested specimen at T = 573 K and r = 50 MPa. Slip traces

are indicated by the arrows. The tensile axis was horizontal

Table 1 Summary of creep stress exponent (n), activation energy for creep (Q), and proposed creep deformation mechanisms for Mg and Mg

alloys

Alloy (wt%) T (K) r (MPa) n Q (kJ/mol) Microstructural feature Proposed mechanism References

Mg 473–600 8.2–21 5.2–6.5 135 ± 10 Extensive basal slip traces Dislocation climb [28]

750–820 3.7–6.3 6 (140 ± 10)

? 295/rExtensive non-basal slip

traces

Cross-slip

750–820 \2.5 1 139 Absence of any visible slip

traces

Nabarro-Herring

diffusion

400–800 [40 7.5–10 92–119 – Non-conservative

motion of jogs on

screw dislocations

gliding in basal slip

planes

[33]

\30 7.5–18.8 160–240 – Cross-slip or non-basal slip(pyramidal slip)

423–523 20–50 5.9 106 – Dislocation climb [34]

6272 J Mater Sci (2012) 47:6263–6275

123

Table 1 continued

Alloy (wt%) T (K) r (MPa) n Q (kJ/mol) Microstructural feature Proposed mechanism References

Mg–0.8Al (solid–

solution)

473–600 \10–20 3 140 ± 10 Extensive basal slip traces and

substructure consists of

random distributed

dislocations

Viscous glide [29]

[10–20 6 Substructure consists of well-

defined subgrains

Dislocation climb

750–810 \10 4 230 ± 15 Extensive non-basal slip traces

and substructure consists of

some isolated but well-

defined sub-boundaries

Cross-slip

AZ91D (as-cast) 393–453 40–115 11 94–220 Existence of non-basal

dislocation segments

Cross-slip [37, 38]

300 60–120 4.6 – – Dislocation climb [36]

AZ31 473–523 – 7 94–102 –

AZ61 Dislocation climb [35]

AZ91 (annealed) 573–623 5 126–132

Mg–5Al–1Sr

(as-cast)

448 50–70 2.4 – Microstructural instability in

the grain boundary region

– [39]

70–85 15.1 – Formation of cavities

423–448 70 – 22 Creep-induced precipitation

448–473 233 – Cross-slip

Mg–5.4Zn–0.6Zr

(powder

metallurgy)

423 20–40 2.2 – – Grain boundary sliding [40, 41]

40–83 6.8 Dislocation glide

373–423 30 – 66 – Grain boundary sliding

423–473 131 Dislocation glide

Mg–(0.7–8.3)Y

(heat-treated)

550–650 \70–150 5 (230 ± 30)–

(290 ± 10)

Basal \a[ slip dominated and

non-basal slip activated

Dislocation climb [42, 43]

[70–150 12 – Power low breakdown

Mg–(0.7–3.9)Y

(solution-treated)

4–40 5–6 Non-basal slip dominated Cross-slip

Mg–8.3Y (aged) 20–60 4 – Viscous glide

Mg–9Y (peak-

aged)

523–573 50–100 5.2 240 ± 20 – Dislocation climb [44]

WE43 (peak-aged) 423–473 200–300 10 118.7 – Dislocation climb [45]

473–523 30–200 4–5 232.9 Dynamic coarsening of

precipitates

573 30–70 5.9 – – – [24]

WE54 (peak-aged) 523–573 50–100 5.1 220 ± 20 Dislocation climb [44]

Mg–(1–5)Y–Nd–

Zn–Zr (as-cast)

523–623 40 – 165–212 – Dislocation climb [46]

Mg–3Y–2Nd–

1Zn–1Mn (as

squeeze casted)

573 30–70 5.9 – – Dislocation climb and

cross-slip[24]

Mg–4Y–3Sm–

0.5Zr (as-

extruded)

473 140–220 4.2 – – Dislocation climb [47]

453–493 180 – 140.6

Mg–4Y–1Sc–1Mn

(squeeze casted,

peak-aged)

523–598 40 – 351–392 – – [10]

573 30–50 8 – –

Mg–(5–10)Gd–

1Sc–1Mn

(squeeze casted,

peak-aged)

548–598 40 – 277 –

573 30–50 7.5 – –

J Mater Sci (2012) 47:6263–6275 6273

123

and r = 50–80 MPa, and was comparable to the

HZ32-T5 at T = 573 K and r = 30 MPa.

(2) At temperatures between 523 and 598 K and stresses

between 30 and 140 MPa, the creep exponent was

close to 5 implying that dislocation creep was the

dominant creep mechanism. The measured average

activation energy (233 ± 18 kJ/mol) was higher than

that for lattice self-diffusion of Mg (135 kJ/mol). This

was considered to be associated with the cross-slip

which was supported by the surface deformation

observations which suggested that non-basal slip and

cross-slip were active at 573 K. Grain boundary

offsets were observed along with grain boundary

cracking suggesting grain boundary sliding was also

an active creep mechanism.

(3) Intergranular creep fracture was the dominant creep–

fracture mode. This is consistent with the calculated

creep damage tolerance parameter k range of 1.2–2.3.

The minimum creep rate and fracture time values fit

the original and MMGR models.

Acknowledgements This work was supported by the National

Natural Science Foundation of China (No. 51074106 and No.

50971089), the Key Hi-Tech Research and Development Program of

China (2009AA033501), the National Key Technology R & D Pro-

gram of China (2011BAE22B01-5), and the International Coop-

eration Fund of Shanghai Science and Technology Committee,

Shanghai/Rhone-Alpes Science and Technology cooperation fund

(No. 06SR07104).

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Table 1 continued

Alloy (wt%) T (K) r (MPa) n Q (kJ/mol) Microstructural feature Proposed mechanism References

Mg–(3–15)Gd–

(2–8)Y (hot-

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523–573 50–100 3.7–4.5 (160 ± 20)–

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Mg–10Gd–3Y–

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Mg–(10–19)Sc

(squeeze casted,

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498–648 40 – 216–262 – – [10]

573 20–45 5–6.5 – –

6274 J Mater Sci (2012) 47:6263–6275

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