Mechanical Properties of A Cu-Based Bulk Metallic Glass ... · slightly increasing trend with Si...

JSC Jang et al (O36), BMG VI, May 11-15, 2008, Xi’an, China

Mechanical Properties of A Cu-Based Bulk Metallic Glass

Microalloyed with Silicon

C. C. Fu1, Y. C. Huang1, I. S. Lee1, P. H. Tsai1, J. S. C. Jang1, L. J. Chang1

1. Department of Materials & Engineering, I-Shou University, Kaohsiung, Taiwan, 840

2. Institute of Materials Science and Engineering, Center for Nanoscience and Nanotechnology, National Sun Yat-Sen

University, Kaohsiung, Taiwan 804

3. Institute of Materials Science and Technology, National Taiwan University of Science and Technology, Taipei,

Taiwan 10607

Abstract

The (Cu42Zr42Al8Ag8)100-xSix amorphous alloy rods, x =0 to 1, with 2 - 6 mm in diameter were

prepared by Cu-mold drop casting method. The thermal properties, microstructure evolution, and

mechanical properties were studied by means of differential scanning calorimetry (DSC), X-ray

diffractometry (XRD), transmission electron microscopy (TEM), hardness test, and compression

test. The XRD result reveals that all of these as-quenched (Cu42Zr42Al8Ag8)100-xSix alloy rods

exhibit a broaden diffraction pattern of amorphous phase. The (Cu42Zr42Al8Ag8)99.5Si0.5 alloy was

found to posses the highest glass forming ability (GFA) as well as the best thermal stability among

the (Cu42Zr42Al8Ag8)100-xSix alloy system. In addition, both of the hardness and yield strength

exhibits a slightly increasing trend with the microalloyed Si content. However, the best mechanical

performance with 2000MPa fracture strength and 3.5 % plastic strain among these

(Cu42Zr42Al8Ag8)100-xSix alloys occurs at the (Cu42Zr42Al8Ag8)99.5Si0.5 amorphous alloy.

Key words: bulk metallic glasses, microalloying, mechanical property, thermal stability

Corresponding author: J. S. C. Jang, e-mail: [email protected]

Introduction


Since the La-base bulk metallic glass (BMG) was first synthesized by copper mold casting

method in 1989 [1], a number of bulk metallic glasses (BMGs) with a wide supercooled liquid

region before crystallization have been prepared in the multicomponent systems such as Pd-[2],

La-[3], Zr- [4,5], Cu-[6-9], Ti-[10], Fe-[11,12], Co-[13], Ni-[14,15], Pt-[16], and Au-[17] based

alloys. Among these BMGs, Cu-base BMGs show some important advantages in engineering

application, including high hardness and strength, large plasticity, high thermal stability and

relatively lower cost [18,19]. Recently, a new Cu-Zr-based BMG with composition of

Cu42Zr42Ag8Al8 [20] has been reported to posses a high glass forming ability (GFA) and can be

produced a BMG rod with diameter of 10 mm by copper mold casting easily. The high GFA of this

alloy system was suggested to result from the formation of highly stabilized liquid in the

Cu-Zr-Ag-Al alloy system. Additionally, the microalloying with metalloid elements has been used

for improving the GFA and thermal stability for many BMGs [21–31]. This method provides an

increase in packing density and so as to increase the GFA, the thermal stability and the mechanical

strength of an amorphous phase. Experimental evidence in our previous study also indicates that

alloying of small metalloid atoms, B or Si, presents the most effective in raising GFA as well as

thermal stability [32–36]. In present study, Cu42Zr42Al8Ag8 alloy which posses high GFA (Trg = 0.63,

Trg = Tg/Tl [37]; γ = 0.382, lg

x

TTT

γ +

= [38]; γm = 0.698, l

gxm T

TT

−=

2γ [39]) was selected as

the base alloy for investigating the effect of microalloying with Si (small atom) on its thermal

properties as well as its mechanical properties.

Experimental Procedures

The pre-alloyed ingots based on the composition of (Cu42Zr42Al8Ag8)100-xSix (x = 0 ~ 1) were

prepared by arc melting of the appropriate mixture of pure elements, such as copper (99.99 wt%

purity), zirconium (99.8 wt% purity), aluminum (99.99 wt% purity), silver (99.99 wt% purity), and

silicon (99.99 % purity), under a Ti-gettered argon atmosphere.. Then the alloy ingots were


remelted in an arc furnace under a purified argon atmosphere. After complete melting, the liquid

alloy was drop cast into the water-cooled Cu mold to form alloy rods with diameter of 2 ~ 6 mm, as

shown in Figure 1. The thermal properties of the as-quenched samples were characterized by TA

Instruments DSC 2920 differential scanning calorimeter (DSC) and high temperature differential

scanning calorimeter (Netsch HTDSC). The as quenched and the annealed structure were examined

by the X-ray diffraction (Scintag X-400 X-ray diffractormeter) with monochromatic Cu-Kα

radiation and transmission electron microscopy (Philip Tenai G2 TEM) with 200 kV. The hardness

was examined by using Akashi MVK-H11 micro hardness tester. The compression strength was

tested by means of MTS 810 mechanical testing system.

Results and Discussion

The results of X-ray diffraction for the (Cu42Zr42Al8Ag8)100-xSix (x = 0 ~ 1) alloy rods with

3mm in diameter shows a broaden maximum in the range of 30°- 50° for all of these alloys in this

study. This indicates that amorphous state of these alloys had been achieved by Cu-mold drop

casting method. A typical outlook of the amorphous rods from 2 to 6 mm without any porosity for

the (Cu42Zr42Al8Ag8)99.5Si0.5 alloy (as shown in Fig. 2) shows the characteristic shining surface. In

addition, the TEM observation also revealed that a uniform amorphous morphology in the as-cast

(Cu42Zr42Al8Ag8)99.5Si0.5 BMG rod with diameter of 3 mm, as shown in Figure 3.

All of the DSC scans of the (Cu42Zr42Al8Ag8)100-xSix (x = 0 ~ 1) alloys presents a clear glass

transition followed by a supercooled liquid region and then exothermic reaction due to

crystallization, as shown in Figure 4. There is only one single peak was observed in the

crystallization for all of these alloys in their DSC curves. It is suggested that the minor silicon

additions do not change the single stage crystallization in this alloy system. In addition, the lowest

liquidus temperature (about 1123 K) occurs at the (Cu42Zr42Al8Ag8)99.5Si0.5 alloy, as show in Table 1.

According to the analyses of Turnbull [40], the best metallic glass forming alloys are at or near deep

eutectic composition and also result in obtaining highest reduced glass transition temperature Trg.


This implies that the alloy (Cu42Zr42Al8Ag8)99.5Si0.5 may be the optimum composition to perform

the best GFA in the (Cu42Zr42Al8Ag8)100-xSix alloy system.

According to the result of DSC analysis as listed in Table 1, the silicon content exhibits slight

effect on the value of ΔTx as well as the values of Trg , γ, and γm. The optimum combination of ΔTx,

Trg γ, and γm values occurs at the alloy composition of (Cu42Zr42Al8Ag8)99.5Si0.5, they are 50 K, 0.63,

0.414 and 0.719, respectively. This suggests that the (Cu42Zr42Al8Ag8)99.5Si0.5 alloy would have the

highest GFA in this alloy system.

By using the Johnson-Mehl-Avrami (JMA) [41] isothermal analysis for volume fraction x

transformed as a function of time t based on the following equation (1):

][-(kt)- x(t) nexp1= (1)

(Cu42Zr42Al8Ag8)100-xSix amorphous alloys were annealed isothermally at several temperatures

between Tg and Tx. To construct the JMA plots, the volume fraction of crystallization at time t was

assumed to be the same as that of heat released. Therefore, the volume crystallization fraction x

which obtained by measuring the partial area under peak up to time t versus the annealing time is

plotted as shown in Figure 5. In parallel, the incubation time as a function of isothermal temperature,

as shown in Figure 6, shows an increasing trend with Si addition. This is an evidence to agree with

the previous reports [35,36], Si would present a positive effect on improving the thermal stability of

some BMGs.

The result of hardness test for these (Cu42Zr42Al8Ag8)100-xSix amorphous alloys exhibits a

slightly increasing trend with Si addition and saturated at the value about 625 in Hv for the

(Cu42Zr42Al8Ag8)99.25Si0.75 amorphous alloy as shown in Fig. 7. This is presumed to be caused from

an increase in the packing density by adding the smaller Si atom in the Cu42Zr42Al8Ag8 alloy system

[28].On the other hand, the results of compression test exhibits similar increasing trend with the Si

content, the higher Si content the higher yield strength, as shown in Fig. 8. However, the optimum

performance with fracture strength more than 2000 MPa and 3.5 % plastic strain occurs at the


(Cu42Zr42Al8Ag8)99.5Si0.5 amorphous alloy.

The entire fracture surface of these (Cu42Zr42Al8Ag8)100-xSix BMG rods after compression test

exhibits more or less area of typical vein pattern, as shown in Fig. 9. However, the fracture surface

of the BMG rods with different Si content presents quite different morphology of vein pattern, such

as the area fraction of vein pattern and the width of vein size. The largest width of vein size around

15 μm was found from the fracture surface of the (Cu42Zr42Al8Ag8)99.5Si0.5 BMG rod, which

performs the optimum combination of fracture strength and plastic strain. This is in agreement with

the suggestion by Inoue [42]. The increase in the diameter of veins suggests the increase in

thickness of the shear deformation region which results in increasing the energy required for plastic

deformation and the final fracture.

Conclusion

Base on the results of thermal analyses, X-ray diffraction, TEM, hardness test, and compression

test for the (Cu42Zr42Al8Ag8)100Si (x = 0 ~ 1) amorphous alloys, the effect of microalloying with Si

on the glass forming ability, thermal properties, and mechanical properties can be summarized as

follows:

(1) The addition of silicon content exhibits a slight effect on increasing the GFA of the

(Cu42Zr42Al8Ag8)100-xSix alloy system. The optimum combination of ΔTx, Trg, γ, and γm values

occurs at the composition of (Cu42Zr42Al8Ag8)99.5Si0.5, they are 50 K, 0.63, 0.414 and 0.719,

respectively. This suggests that the (Cu42Zr42Al8Ag8)99.5Si0.5 alloy would have the highest GFA

in this alloy system.

(2) The hardness of these (Cu42Zr42Al8Ag8)100-xSix BMGs presents a slightly increasing trend with

Si addition and saturates at the value about Hv 625 for the (Cu42Zr42Al8Ag8)99.25Si0.75

amorphous alloy.

(3) The result of compression test shows that the yield strength increases with Si addition.

However, the optimum performance with fracture strength more than 2000 MPa and 3.5 %


plastic strain occurs at the (Cu42Zr42Al8Ag8)99.5Si0.5 BMG.

(4) The largest width of vein size around 15 μm was found from the fracture surface of the

(Cu42Zr42Al8Ag8)99.5Si0.5 BMG rod, which performs the optimum combination of fracture

strength and plastic strain. The increase in the diameter of veins suggests the increase in

thickness of the shear deformation region which results in increasing the energy required for

plastic deformation and the final fracture.

Acknowledgement

The authors would like to gratefully acknowledge the sponsorship from the National Science

Council of ROC under the project NSC 96-2218-E-110-001. In addition, the authors are also very

grateful for the assistance of X-ray diffraction and TEM by the Micro and Nano Laboratory,

Department of Materials Science and Engineering, I-Shou University.

References

1. A. Inoue, T. Zhang, and T. Masumoto, Mater. Trans. JIM, 30 (1989) 965.

2. A. Inoue, N. Nishiyama, and T. Matsuda, Mater. Trans. JIM 37 (1996) 181.

3. Inoue, H. Yamaguchi, T. Zhang, and T. Masumoto, Mater. Trans. JIM. 31 (1990) 104.

4. T. Zhang, A. Inoue, and T. Masumoto, Mater. Trans. JIM 32 (1991) 1005.

5. A. Pecker, W. L. Johnson, Appl. Phys. Lett., 63 (1993) 2342.

6. Inoue,W. Zhang, T. Zhang, and K. Kurosaka, Acta Mater. 49 (2001) 2645.

7. Inoue,W. Zhang, T. Zhang, and K. Kurosaka, J Non-Cryst. Solids 304 (2002) 200.

8. J. Eckert, J. Das, K. B. Kim, F. Baier, M. B. Tang, W. H. Wang, and Z. F. Zhang,

Intermetallics 14 (2006) 876.

9. W. Zhang, F. Jia, Q. Zhang, and A. Inoue, Mater. Sci. Eng., A459 (2007) 330.

10. T. Zhang, A. Inoue, and T. Masumoto, Mater. Trans. JIM 39 (1998) 1001.

11. J. P. Lu, C. T. Liu, J. R. Thompson, and W. D. Porter, Phts. Rev. Lett.92 (2004) 245503.


12. J. Shen, Q. J. Chen, J. F. Sun, H. B. Fan, and G. Wang, Appl. Phys. Lett. 86 (2005) 151907.

13. T. Itoi, A. Inoue, Mater. Trans. JIM 41 (2000) 1256.

14. S. Yi, T. G. Park, and D. H. Kim, J. Mater. Res. 15 (2000) 2425.

15. T. Zhang, A. Inoue, Mater. Trans. JIM 43 (2002) 708.

16. J. Schroers, W. L. Johnson, Appl. Phys. Lett., 84 (2004) 3666.

17. J. Schroers, B. Lohwongwatana, W. L. Johnson, and A. Peker, Appl. Phys. Lett., 87 (2005)

061912-1.

18. A. Inoue, W. Zhang, T. Zhang, and K. Kurosaka, Mater. Trans. 42 (2001) 1149.

19. T. Zhang, A. Inoue, Mater. Trans. 43 (2002) 1367.

20. Q. Zhang, W. Zhang, and A. Inoue, Scri. Mater., 55 (2006) 711.

21. W. H.Wang, Z. Bian, P. Wen, M. X. Pan, and D. Q. Zhao, Intermetallics, 10 (2002) 1249.

22. Z. P. Lu, C. T. Liu, J. Mater. Sci., 39 (2004) 3965.

23. C. T. Liu, Z. P. Lu, Intermetallics, 13 (2005) 415.

24. A. Gebert, J. Eckert, and L. Schultz, Acta Mater., 46 (1998) 5475.

25. W. H. Wang, Q. Wei, and H. Y. Bai, Appl. Phys. Lett., 71 (1997) 58.

26. H. Choi-Yim, R. Busch , and W. L. Johnson. J Appl. Phys. 83 (1998) 7993.

27. C. Ma, H. Soejima, K. Amiya, N. Nishiyama, and A. Inoue, Mater. Trans, 45 (2004) 3223.

28. W. H. Wang, Prog. In Mater. Sci., 52 (2007) 540.

29. Z. P. Lu, C. T. Liu, Acta Mater., 50 (2002) 3501

30. Z. P. Lu, C. T. Liu, W. D. Porter, Appl. Phys. Lett. 83 (2003) 2581.

31. C. T. Liu, M. F. Chisholm, and M. K. Miller. Intermetallics 10 (2002) 1105

32. J. S. C. Jang, L. J. Chang, Y. T. Jiang, and P.W. Wong, Mater. Sci. Forum, 426–432 (2003)

1879.

33. J. S. C. Jang, Y. W. Chen, L. J. Chang, and G. J. Chen, Mater. Chem. Phys. 88 (2004) 227.

34. J. S. C. Jang, Y. C. Huang, C. H. Lee, I. S. Lee, and L. J. Chang, Mater. Sci. Forum, 561-565

(2007) 1341.


35. J. S. C. Jang, S. F. Tsao, L. J. Chang, G. J. Chen, and J. C. Huang, J. Non-Cryst. Solids 352

(2006) 71.

36. J. S. C. Jang, L. J. Chang, T. H. Hung, J. C. Huang, and C. T. Liu, Intermetallics 14 (2006)

951.

37. W. L. Johnson, MRS Bull. 24 (10) (1999) 42.

38. Z. P. Lu, C. T. Liu, Intermetallics 12 (2004) 1035.

39. X. H. Du, J. C. Huang, C. T. Liu, and Z. P. Lu, Jap. J. Appl. Phys., 101 (2007) 086108.

40. D. Turnbull, Contemp. Phys. 10 (1969) 473.

41. M. Avrami : J. Chem. Phys.,7 (1939)1103-1112.

42. A. Inoue, Bulk Amorphous Alloys - Practical Characteristics and Applications, Trans Tech

Publ. Ltd, Switzerland (1999), pp. 5.

Table caption:


Table 1.Thermal parameters of the (Cu42Zr42Al8Ag8)100-xSix (x = 0 ~ 1) alloys

Figure captions:

Fig. 1. XRD patterns of as-quenched (Cu42Zr42Al8Ag8)100-xSix (x = 0 ~ 1) BMG rods with 3 mm in

diameter.

Fig. 2. As-cast (Cu42Zr42Al8Ag8)99.5Si0.5 BMG rods with diameter from 2 to 6 mm

Fig. 3. HRTEM image of the as-cast (Cu42Zr42Al8Ag8)99.5Si0.5 BMG rod with diameter of 3 mm

Fig. 4. DSC plots of (Cu42Zr42Al8Ag8)100-xSix (x = 0 ~ 1) BMG rods

Fig. 5. The fraction transformed versus the annealing time of isothermal DSC plots

for(Cu42Zr42Al8Ag8)100Si (x = 0 ~ 1) BMG rods.

Fig. 6. Incubation time as a function of isothermal temperature for(Cu42Zr42Al8Ag8)100Si (x = 0 ~ 1)

BMG rods.

Fig. 7. Hardness as a function of Si addition for (Cu42Zr42Al8Ag8)100Six (x = 0 ~ 1) BMG rods

Fig. 8. Compression stress-strain curves of the (Cu42Zr42Al8Ag8)100-xSix BMG rods with diameter of

2 mm

Fig. 9. SEM images of the fracture surface after compression test for the (Cu42Zr42Al8Ag8)100-xSix

BMG rods with diameter of 2 mm


Table 1. Thermal parameters of the (Cu42Zr42Al8Ag8)100-xSix (x = 0 ~ 1) alloys

Si content

(at%) Tg (K) Tx (K) T△ x (K) Tm (K) Tl (K) T△ l (K) Trg γ γm

X=0 703 748 45 1098 1136 38 0.62 0.407 0.698

X=0.25 703 755 52 1094 1131 37 0.62 0.412 0.713

X=0.5 708 758 50 1089 1123 34 0.63 0.414 0.719

X=0.75 715 761 47 1100 1126 25 0.63 0.413 0.716

X=1 715 762 47 1097 1128 31 0.63 0.413 0.717

△Tx= Tx –Tg, △Tl = Tl –Tm,


20 30 40 50 60 70 80

x=1

x=0.25

x=0.5

x=0.75

x=0

Inte

nsity

(arb

. uni

ts)

2 θ

Fig. 1 XRD patterns of as-quenched (Cu42Zr42Al8Ag8)100-xSix (x = 0 ~ 1) BMG rods with 3 mm in diameter.


Fig. 2. As-cast (Cu42Zr42Al8Ag8)99.5Si0.5 BMG rods with diameter from 2 to 6 mm


Fig. 3. HRTEM image of the as-cast (Cu42Zr42Al8Ag8)99.5Si0.5 BMG rod

with diameter of 3 mm

5 nm


400 500 600 700 800 900

TX

X=0.75

X=0.25

X=0.5

X=0

Temperature (K)

Hea

t flo

w (e

xo.)

X=1

Tg

Fig. 4 DSC plots of (Cu42Zr42Al8Ag8)100-xSix (x = 0 ~ 1) BMG rods


0 500 1000 1500 2000 25000

50100

753 K748 K

x=0.75

x=0.5

x=0.25

x=0

Annealing time (s)

x=1

743 K

050

100

050

100

Cry

stal

lizat

ion

frac

tion

(%)

050

100

050

100

725 K730 K735 K

730 K735 K740 K745 K

733 K738 K743 K

745 K750 K755 K

Fig. 5 The fraction transformed versus the annealing time of isothermal DSC plots

for(Cu42Zr42Al8Ag8)100Si (x = 0 ~ 1) BMG rods.


725 730 735 740 745 750 7550

200

400

600

800

1000

1200

1400

Incu

batio

n tim

e (s

)

Temperature (K)

x=0 x=0.25 x=0.5 x=0.75 x=1

Fig. 6 Incubation time as a function of isothermal temperature for(Cu42Zr42Al8Ag8)100Si (x = 0 ~ 1)

BMG rods.


0.0 0.2 0.4 0.6 0.8 1.0450

500

550

600

650

700

750

Har

dnes

s (H

v)

Si content (at%)

Fig. 7 Hardness as a function of Si addition for (Cu42Zr42Al8Ag8)100Six (x = 0 ~ 1) BMG rods



0.00 0.02 0.04 0.06 0.08 0.10 0.120

500

1000

1500

2000

2500

2 %

Strain

1 Si0.75 Si0.5 Si

0.25 Si

Stre

ss (M

Pa)

0 Si

Strain rate: 5x 10-4 s-1

Fig. 8 Compression stress-strain curves of the (Cu42Zr42Al8Ag8)100-xSix BMG rods



Fig. 9 SEM images of the fracture surface after compression test for the (Cu42Zr42Al8Ag8)100-xSix BMG

rods with diameter of 2 mm

a 0 at% Si b 0.25 at% Si

c 0.5 at% Si d 0.75 at% Si

e 1.0 at% Si

Mechanical Properties of A Cu-Based Bulk Metallic Glass ... · slightly increasing trend with Si...

Documents

Transcript of Mechanical Properties of A Cu-Based Bulk Metallic Glass ... · slightly increasing trend with Si...