IMPROVEMENT OF HOT DUCTILITY IN Al-Mg BASE ALLOYS CAUSED BY SMALL AMOUNTS OF ADDITIONAL ELEMENTS

Keitaro HORIKAWA, Shigeru KURAMOTO and Motohiro KANNO

The University of Tokyo Department of Materials Science, School of Engineering,

7-3-1 Hongo, Bunkyo-ku, Tokyo 113, Japan

Abstract

An Al-5.5mol%Mg alloy containing traces of sodium, calcium or strontium as an impurity shows high temperature embrittlement based on intergranular fracture. In the present study, effect of small amounts of silicon on hot ductility of an Al-Mg alloy containing traces of sodium or strontium was examined at elevated temperatures. High temperature embrittlement caused by 2ppm of sodium was successfully suppressed by the addition of about 1500molppm of silicon, and that by 2ppm of strontium was suppressed by the addition of lOOOppm silicon. The Al-Mg-Na or Al-Mg-Sr alloy showed almost intergranular fracture surface at about 300°C, while the Al-Mg-Na alloy contaimng 1480ppm silicon and the Al-Mg-Sr alloy containing 940ppm silicon a transgranular one. Energy dispersive X-ray analysis revealed that Si-bearing compounds trap strontium in the Al-Mg-Sr-Si alloy. Thus, it is concluded that silicon can scavenge sodium or strontium from grain boundaries resulting in the suppression of high temperature embrittlement.

Automotive Alloys 1999 Edited by S.K. Das

The Minerals, Metals & Materials Society, 2000

181

Introduction

Aluminum alloys containing magnesium between 5 to 10 mass% are now being developed for automotive parts, since they have good formability. However, it has been reported that an A1-5.5mol%Mg alloy shows high temperature embrittlement based on intergranular fracture when tested at 300°C at an initial strain rate of 8.3xlO"V (1). Embrittlement temperature is reported to be shifed to high as strain rate increases (2). This embrittlement is thus closely related to unfavorable hot rolling characteristics of the Al-Mg alloy. The authors reported recently that this embrittlement is caused by a trace amount of sodium of only 0.1 molppm in a coarse-grained Al-5.5mol%Mg alloy (3,4). Further, traces of only 2 molppm of strontium or calcium were shown to bring about high temperature embrittlement resulting from these segregation to grain boundaries (5).

To improve the hot rolling characteristics of the Al-Mg alloy, investigation is needed on the way to eliminate the effect of harmful impurities. Talbot and Ransley reported that the hot working characteristic of an Al-Mg alloy containing 22 massppm of sodium is improved by the addition of 100 massppm of bismuth (6). The authors have also reported that antimony addition of 2 molppm can suppress high temperature embrittlement caused by 2 molppm of sodium (7). A similar effect can be expected by adopting the other additional elements for high temperature embrittlement caused by sodium, strontium or calcium. In this study, silicon was selected as an additive. This is because silicon is assumed to have high affinity with sodium or strontium, based on the fact that morphology of a silicon rich phase in an Al-Si casting alloy is known to be modified by the addition of a small amount of sodium or strontium (8). The aim of this work was to determine whether a small amount of silicon can suppress the high temperature embrittlement of an Al-Mg alloy containing traces of sodium or strontium.

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Experimental Procedure

Seven kinds of Al-5%Mg alloys were melted using a graphite crucible in a high frequency induction furnace under argon atmosphere of 105Pa (Table 1). Aluminum of 99.999mass% purity and magnesium of 99.98mass% purity were used. Trace elements were added to Al-Mg alloys using the master alloys Al-0.06mol%Na, Al-0.94mol%Sr, and Al-12mol%Si prepared in our laboratory. Ingots were homogenized at 430°C for 18hr in argon of 105Pa, cold-swaged by 70% and machined into round tensile specimens having a gage length of 10mm and a diameter of 4mm. Tensile test specimens with silicon added were annealed at 550°C and those without silicon added at 510°C to get a grain size of about 300(xm. Trace elements were analyzed by emission spectrochemical analysis.

Tensile tests were performed at temperatures ranging from room temperature to 450°C at an initial strain rate of 8.3x10"V. Hot ductility was evaluated after testing by measuring the reduction in area. Fracture surfaces were observed using a scanning electron microscope (SEM). Compounds appearing in the silicon added alloys were analyzed with energy dispersive X-ray spectroscopy (EDXS) after mechanical- and electropolishing in an electrolytic solution of nitric acid and methylalcohol.

Table 1 Chemical Composition of Al-Mg Alloys

ELEMENT ALLOY Al-Mg-Na Al-Mg-Na-200ppmSi Al-Mg-Na-700ppmSi Al-Mg-Na- ISOOppmSi Al-Mg-Sr Al-Mg-Sr-200ppmSi Al-Mg-Sr-lOOOppmSi

Mg (mol%)

5.85 5.41 6.73 5.63 5.41 5.63 5.41

Na*

1.2 1.2 1.2 2.2 0.1 0.1 0.1

Sr*

<0.1 <0.1 <0.1 <0.1 2.0 3.0 2.0

Si*

25 70

700 1480

18 230 940

Fe*

<1 2

<1 1 2 <1 1

Al*

bal. bal. bal. bal. bal. bal. bal.

*molppm

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Results

Silicon addition to Al-Mg-Na alloys

The effect of small amounts of silicon on reduction in area versus temperature curves of an Al-Mg-Na alloy is shown in Figure 1. High temperature embrittlement appears in an Al-Mg-Na alloy at around 300°C and the value of reduction in area at the bottom of a ductility trough is as low as 10% as has been reported previously (5). In this case, the alloy exhibited early fracture during uniform deformation. Silicon additions of less than 700ppm to the above alloy had almost no influence on hot ductility. On the other hand, the embrittlement was drastically suppressed and such ductility trough almost disappeared by the addition of 1480ppm silicon and the value of reduction in area at the bottom of the shallow ductility trough was as high as about 70%. This value is similar to that of the antimony bearing Al-Mg-Na alloys as reported elsewhere (7). In contrast to the alloy without silicon, the Al-Mg-Na- 1480ppmSi alloy fractured after local necking through uniform deformation.

p. L i I ■ I . I . | . _

0 100 200 300 400 500

Temperature, T I °C

Figure 1: Effect of Xppm silicon on reduction in area vs. tensile testing temperatures in Al-5.5%Mg-2ppmNa alloys

184

Fracture surfaces of the Al-Mg-Na and the Al-Mg-Na- 1480ppmSi alloy tensile tested at 300°C are indicated in Figure 2. The former alloy shows slight contraction of the cross section corresponding to Figure 1 and a typical intergranular fracture surface, while, the latter shows heavy necking and a mixture of intergranular and transgranular fracture surfaces. Thus, it is clear that silicon addition can suppress early intergranular cracking during deformation.

Figure 2: SEM fractographs of Al-Mg-Na (a) and Al-Mg-Na- 1480ppm Si (b) specimens tested at 300°C.

Silicon addition to Al-Mg-Sr alloys

Figure 3 shows the effect of silicon on hot ductility versus temperature curves of Al-Mg-Sr alloys. The Al-Mg-Sr alloy also showed high temperature embrittlement but the embrittlement was milder that of the Al-Mg-Na alloy; the ductility trough became narrow and shallow. Silicon addition of 230ppm had no influence on hot ductility of the Al-Mg-Sr alloy. Silicon addition of 940ppm, in contrast, sufficiently suppressed the high temperature embrittlement and the value of reduction in area at the bottom of the ductility trough was as high as about 80%.

Figure 4 shows the fracture surfaces of the Al-Mg-Sr and the Al-Mg-Sr-940ppmSi alloys tested at the embrittlement temperature. It is noted that the fracture mode of the Al-Mg-Sr alloy is intergranular but that 940ppmSi addition changed it into intergranular. Therefore, it is obvious that suppression of high temperature embrittlement caused by strontium is due to the variation of fracture mode from intergranular to transgranular.

185

(0 0) <

c O o

or

100

80

60

40

20

AI-5%Mg 2ppmSr -X ppmSi

- D - X=(18) - S - X=230 _■_ X=940

-L 100 200 300 400 500

Temperature, T I °C

Figure 3: Effect of Xppm silicon on reduction in area vs. tensile testing temperatures in Al-5.5%Mg-2ppmSr alloys

To determine why the embrittlement was suppressed by the addition of 940ppm silicon, electropolished specimens were observed using SEM. Figure 5 demonstrates that compounds distributed uniformly all over the surface of the specimen and the size of a compound was about 2-3 \xm (Figure 5(b)). An analysis by EDXS revealed that a compound consisting of silicon, strontium, and magnesium was formed in the Al-Mg-Sr-Si alloy as shown in Figure 6. Although a peak for aluminum is seen in this figure, this peak may not have been from the compound but from the matrix.

186

Figure 4: SEM fractographs of Al-Mg-Sr (a) and Al-Mg-Sr-940ppm Si (b) specimens tested at 250°C.

Figure 5: SEM images of electrolytically polished surface of the Al-Mg-Sr-940ppmSi specimen ((a): low magnification (b): high magnification)

187

CO o tn (0

■ ^

la <

i ! -' i — i — i — i — r

MKal

Total CoiBtt'49103. Unwr V3-3000 -, 1 , , Sp 1 1 1 |—

Magnified

Si Sr

15M

1001

Energy / keV

Figure 6: EDXS analysis of the compound in the Al-Mg-Sr-940ppmSi alloy in Figure 5

Discussion

It has been reported that antimony of only 2ppm can suppress high temperature embrittlement caused by 2ppm of sodium through the formation of sodium-antimony bearing compounds (7). In addition to this, the present study revealed that the additional silicon of as much as 1480ppm or 940ppm could suppress high temperature embrittlement caused by 2ppm of sodium or strontium, respectively. Optical and electron microscopy revealed that silicon-bearing compounds were densely distributed in the Al-Mg-Sr-940ppmSi alloy, but not in the Al-Mg-Sr-230ppmSi or the Al-Mg-Sr alloy after annealing as shown in Figure 7. Though the solid solubility limit of silicon at about 500°C in an Al-5.5%Mg alloy has not been reported, additional magnesium is known to greatly reduce that limit (9). This suggests that silicon of about 700ppm in an Al-5.5mol%Mg alloy may be in a

188

solid solution state at the annealing temperature. Accordingly, it is reasonable to say that the suppression of high temperature embrittlement by the addition of silicon is related to the formation of these compounds, which is similar to the case of 2ppm antimony addition.

Distribution density of compounds in the Al-Mg-Sr-940ppmSi alloy is much higher than that of sodium-antimony bearing compounds in the Al-Mg-2ppmNa-2ppmSb alloy as reported, however (7). In addition, the intensity of the strontium peak from the silicon-bearing compound is very weak as compared to that of sodium from the sodium-antimony compound in the Al-Mg-2ppmNa-2ppmSb alloy as shown in Figure 6. Therefore, the mechanism of the suppression of the embrittlement by the addition of silicon is apparently different from that by antimony addition to an Al-Mg-Na alloy. It is thus reasonable to conclude that sodium or strontium may be trapped preferentially at the interface between silicon-bearing compounds and the matrix in the Al-Mg-Na- 1480ppmSi or the Al-Mg-Sr-940ppmSi alloy during annealing.

WTiQOQ

q Figure 7: Optical microscopic images of the Al-Mg-Sr (a), the Al-Mg-Sr-230ppmSr (b),

and the Al-Mg-Sr-940ppmSr alloy (c)

189

Conclusions

The effect of small amounts of silicon on hot ductility of coarse-grained Al-5.5mol%Mg alloys containing 2molppm of sodium or strontium was examined at temperatures ranging from ambient temperature to 450°C. A series of experiments has shown that the high temperature embrittlement caused by traces of either element is almost completely suppressed by the addition of silicon at 1480ppm or 940ppm, respectively.

The mechanism of suppression of high temperature embrittlement by the addition of silicon may be different from the case of antimony addition to an Al-Mg-Na alloy. It is supposed that sodium or strontium is trapped at the interface between silicon-bearing compounds and the matrix in an Al-Mg-Na-Si or an Al-Mg-Sr-Si alloy.

Acknowledgments

This work was supported by the Light Metal Educational Foundation Inc. (Osaka, Japan) The authors wish to thank Furukawa Electric Co., Ltd. for analyses of bulk materials. They are also grateful to Mr. K. Ueda and Mr. H. Koyo for their experimental work.

References

1. C. E. Ransley and D.E. J. Talbot, J. Inst. Metals. 88, (1959-60) 150-158. 2. M. Otsuka and R. Horiuchi, J. Jpn. Inst. Metals, 49, (1984) 688. 3. H. Okada and M. Kanno, Scripta Materialia, 37(6)(1997), 781-786. 4. K. Horikawa, S. Kuramoto and M. Kanno, J. Jpn. Inst. Light Metals, 48(8)(1998), 371-

374. 5. K. Horikawa, S. Kuramoto and M. Kanno, Scripta Materialia. 39,7(1998), 861-866. 6. E. J. Talbot and D. A. Granger, J. Minerals, Metals and Materials Society. 47, (1995),

44-46. 7. K. Ueda, K. Horikawa and M. Kanno, Scripta Materialia, 37(8)(1997), 1105-1110. 8. J. E. Hatch, ed., Aluminum, (American Society for Metals, 1984). 9. Mondolfo, ed., Aluminum Alloys, (Butterworth & co. Ltd., 1976).

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