Hiroyuki Yamada
*, Keitaro Horikawa and Hidetoshi Kobayashi
Department of Mechanical Science and Bioengineering, Graduate School of Engineering Science, Osaka University, Toyonaka 560-8531, Japan
In order to elucidate the mechanism of the high temperature embrittlement of Al-5 mass% Mg alloys containing traces of sodium, microscopic observation of grain boundaries during quasi-static and impact tensile deformation was carried out. The high temperature embrittlement during the quasi-static deformation was caused by a combination of grain boundary sliding and grain boundary serration. On the other hand, the high temperature embrittlement during the impact deformation was caused by the localization of stress concentration at the grain
boundaries. This localization was due to planar slips. [doi:10.2320/matertrans.L-M2009824]
(Received April 9, 2009; Accepted July 31, 2009; Published November 11, 2009)
Keywords: aluminum-magnesium alloy, high temperature embrittlement, sodium, strain rate, grain boundary microstructure
1. Introduction
In recent years, Al-Mg alloys have been widely used because of their high corrosion resistance, high strength, and high ductility at room temperature. However, it is known1–4)
that coarse-grained Al-Mg alloys containing 5 mass% Mg (hereafter, mass% is referred to as %) show low ductility around 573 K, as indicated by the grain boundary (GB) fracture in the quasi-static tensile test. This loss of ductility is called high temperature embrittlement (HTE). The HTE of Al-5% Mg alloys is believed to be caused by the segregation of trace sodium below 1 mol ppm (hereafter, mol ppm is referred to as ppm) to GBs and the lowering of the GB strength.2–6)On the other hand, Lynch7)suggested that HTE
is due to the formation of low melting point sodium-rich phases or particles rather than the segregation of trace sodium at the GBs of Al-5% Mg alloys in comparison with the case of Al-Li alloys. However, sodium segregation has not been detected on GB fracture surface by means of Auger electron spectroscopy.2)
It has been reported1) that the temperature at which the
ductility minimum occurred increased and the minimum value of the reduction in area was unchanged when the strain rate in the quasi-static test increased. The dependence of the intermediate strain rate (105–101s1) on the HTE of Al-5% Mg alloys has been reported,1)but the effect of a high strain rate (approximately 102s1) on the HTE was not clarified in detail. The authors8)have shown that the HTE of
the Al-5% Mg alloy containing 2.0 ppm of sodium decreased with increasing testing temperature and that no recovery of ductility above 673 K was observed when the specimen was tested at an average strain rate of5:0102s1.
Otsuka and Horiuchi9)have reported that the HTE of an
Al-5% Mg binary alloy was caused by an intergranular fracture resulting from the growth and coalescence of cavities at GBs in the quasi-static deformation. However, the mechanism of HTE in the Al-Mg alloy has not been clarified under the same conditions as those in which the effect of traces of sodium was considered in the previous study.9)
Lynch7)also pointed out that the characterization of the effect
of traces of sodium on GB microstructures, such as GB sliding, migration, and cavity formation, will help understand the mechanism of HTE in Al-Mg alloys. The purpose of this study is to elucidate the mechanism of HTE in Al-5% Mg alloys that contain traces of sodium. The elucidation is based on the microscopic observation of GBs during tensile deformation.
2. Experimental Procedure
Two kinds of Al-5% Mg alloys were melted and cast in an argon atmosphere of 0.1 MPa using Al of 99.999% purity (sodium concentration<0.5 ppm) and magnesium of 99.98% purity (sodium concentration <0.9 ppm), and the Al-5% Mg alloy containing 2.0 ppm of sodium (Al-5% Mg-2 ppm Na alloy) and that containing 0.10 ppm of sodium (Al-5% Mg-0.1 ppm Na alloy) were prepared. The chemical compositions are shown in Table 1. The alloy ingots were homogenized at 703 K for 18 h in the Ar atmosphere and subsequently cold-rolled by 70%. Tensile test pieces with a gage length of 10.0 mm (ls), a gage width of 5.0 mm, a thickness of 2.0 mm, and a fillet radius of 1.0 mm were machined from the cold-rolled sheets. All the test pieces were then annealed at 783 K for 0.5 h in an air atmosphere to obtain an average grain size of 300mm.
The quasi-static test was carried out at a strain rate of
8:3104s1. The impact test was performed by using the split Hopkinson pressure bar (SHPB) method. Dynamic stress-strain curves can be obtained by incorporating elemen-tary one-dimensional elastic wave propagation theory10–13)
[image:1.595.304.550.713.776.2]into the SHPB method. Figure 1 shows the SHPB apparatus
Table 1 Chemical compositions of Al-5% Mg alloys (mol ppm).
Alloy Mg Na Si Fe Al
Al-5% Mg
-2 ppm Na alloy 4.9 2.0 16.4 — Bal.
Al-5% Mg
-0.1 ppm Na alloy 5.1 0.10 8.4 2.2 Bal.
mass%
for the impact tensile test. An incident bar and a transmitted bar (S45C) were both 20 mm in diameter. When a striker tube (A5052) impacts a yoke, a tensile stress wave propagates through the input bar. When the tensile stress wave reaches the specimen, part of the stress wave is reflected at the specimen, while remaining part of the stress wave propagates through the output bar. The stress waves can be measured by semiconductor strain gages attached on the input and the output bars. By applying the elementary one-dimensional elastic wave propagation theory, we can determine the nominal stressðtÞ, the nominal strain"ðtÞand the strain rate
_
"
"in the specimen as follows:
ðtÞ ¼AE
As
"tðtÞ ð1Þ
"ðtÞ ¼2c0
ls
Zt
0
½"iðtÞ "tðtÞdt ð2Þ
_
" "¼2c0
ls
½"iðtÞ "tðtÞ ð3Þ
whereAsis the cross-sectional area of the specimen,Ais the cross-sectional area of the elastic bars,Eis Young’s modulus of the input and output bars,c0is the velocity of elastic wave in the input and output bars (c0¼
ffiffiffiffiffiffiffiffi
E=
p ¼
5250m/s, : density of the input and output bars) and"iðtÞand"tðtÞare the
incident and transmitted waves, respectively. In the impact test, an average strain rate was 5:0102s1, which was calculated by eq. (3).
Both the quasi-static test and the impact test were carried out at temperatures ranging from room temperature to 773 K using an infrared image furnace. The hot ductility was evaluated by measuring the fracture strain. The fracture surfaces and GB microstructures were observed using a scanning electron microscope (SEM). The amount of GB sliding was measured at temperatures ranging from 473 K to 773 K in the same manner as that reported9)by Otsuka and Horiuchi. Before performing the tensile test, more than 20 marker lines parallel to the stress axis were made on the polished surface at line intervals of 200mm. The displace-ment component of GB sliding perpendicular to the stress axis, w, was measured on the specimen surface after the tensile test. The strain due to GB sliding ("gbs) can be
evaluated14–17)as follows:
"gbs ¼kww=dd ð4Þ
wherekis a geometric constant, which is equal to 1.5 for the displacement perpendicular to the stress axis, ww is the average value of w, and dd is the average grain size. The amount of GB sliding was evaluated mainly by calculating the ratio of "gbs to the fracture strain ("F). In order to
investigate the GB morphology in detail, the specimen strained by the quasi-static test was fractured by coating
liquid gallium (melting point: 303 K) directly onto the specimen surface.18,19)
3. Results and Discussion
Figure 2 shows the effects of the strain rate and traces of sodium on hot ductility in the Al-5% Mg alloys. As expected, HTE appeared around 573 K in the quasi-static test for the Al-5% Mg-2 ppm Na alloy. The minimum value of the fracture strain was approximately 4% at 573 K. In both the quasi-static test and the impact test at temperatures ranging from 473 to 773 K, the fracture strain increased when the amount of sodium was decreased. However, as previously reported,8) the behavior of the hot ductility was different in the testing mode. In the impact test, the hot ductility decreased with increasing testing temperature, and no recovery of the ductility above 673 K was observed. A slight improvement in the ductility was observed at 673 K in the impact test when the amount of sodium was 0.1 ppm. However, a marked decrease in the ductility appeared in both the Al-5% 0.1 ppm Na alloy and the Al-5% Mg-2 ppm Na alloy when the impact test was conducted at 773 K. It was also obvious that the temperature lowering the hot ductility shifted to higher temperatures when the impact test was performed.
In order to understand the reason for the change in the fracture strain due to the differences in the sodium amounts or strain rates, the fracture surfaces for the Al-5% Mg-0.1 ppm Na alloy and the Al-5% Mg-2 ppm Na alloy tested at 573 K and 673 K, as shown in Fig. 3, were compared. In the quasi-static tests, the intergranular fracture was observed in
Fig. 1 Split Hopkinson pressure bar apparatus for the impact test.
Fig. 2 Effects of the strain rate and traces of sodium on the hot ductility in
[image:2.595.113.486.71.148.2] [image:2.595.310.541.179.370.2]both alloys tested at 573 K (Figs. 3(a), (e)), where a number of intergranular dimples were observed at the GBs. At 573 K, the average size of the intergranular dimples observed in the Al-5% Mg-0.1 ppm Na alloy was larger than that observed in the Al-5% Mg-2 ppm Na alloy. On the other hand, at 673 K, the transgranular fracture surfaces were identified after performing the quasi-static test for both alloys (Figs. 3(b), (f)). In the impact tests, the fracture surfaces were composed of transgranular dimples and GBs in both alloys tested at 573 K (Figs. 3(c), (g)). The Al-5% Mg-0.1 ppm Na alloy impact-tested at 673 K (Fig. 3(d)) showed a combination of transgranular and intergranular fractures, while the Al-5% Mg-2 ppm Na alloy impact-tested at 673 K showed almost entirely intergranular fractures (Fig. 3(h)), where the ledge patterns that represent the trace of coarse planar slips20,21) were observed. It is noted that the decrease in the sodium content led to a change in the fracture mode from the intergranular fracture to a combination of transgranular and
intergranular fractures even when the impact test was performed at 673 K. Although Zhanget al.22)suggested that
HTE was related to the formation of low melting point sodium-rich liquid phases that precipitates at GBs by means of thermodynamic simulation, few sodium-rich liquid phase distributions were observed at intergranular fracture surfaces in the present study.
Figure 4 shows the effects of strain rate and traces of sodium on the fraction of GB sliding in the Al-5% Mg-0.1 ppm Na alloy and the Al-5% Mg-2 ppm Na alloy. In the quasi-static test, the total amount of GB sliding decreased when the sodium content was decreased, particularly at 573 K and 673 K. Since the amount of GB sliding, which was evaluated by calculating the ratio "gbs="F, was
influ-enced by the fracture strain, the difference in the amount of GB sliding was compared by applying the same strain ("¼0:032 at 573 K) in both alloys. The amount of GB sliding was once again evaluated by calculating the ratio "gbs=0:032 for the Al-5% Mg-0.1 ppm Na alloy at 573 K.
Such comparisons reveal that the ratio "gbs=0:032 for the
Al-5% Mg-0.1 ppm Na alloy at 573 K was 15%, which is approximately half of the corresponding ratio for the Al-5% Mg-2 ppm Na alloy. The Al-5% Mg-0.1 ppm Na alloy strained by 0.032 at 573 K was then fractured by liquid gallium. Figure 5 shows SEM micrographs of the GB
Fig. 3 SEM micrographs of the fracture surfaces in the Al-5% Mg-0.1 ppm
Na alloy—Figs. 3(a)–(d)—and the Al-5% Mg-2 ppm Na alloy— Figs. 3(e)–(h); (a) and (e): static test at 573 K; (b) and (f): quasi-static test at 673 K; (c) and (g): impact test at 573 K; (d) and (h): impact test at 673 K.
Fig. 4 Effects of the strain rate and traces of sodium on the fraction of GB
sliding in the Al-5% Mg-0.1 ppm Na alloy and the Al-5% Mg-2 ppm Na alloy.
Fig. 5 SEM micrographs of the GB surfaces in the Al-5% Mg-0.1 ppm Na
[image:3.595.312.540.71.271.2] [image:3.595.55.283.75.489.2] [image:3.595.311.543.332.419.2]surfaces in the Al-5% Mg-0.1 ppm Na alloy fractured by liquid gallium after the pre-strain (plastic strain = 0.032 at 573 K) and the Al-5% Mg-2 ppm Na alloy fractured by the quasi-static test at 573 K (magnified view of Fig. 3(e)). Very few dimples were observed on the GB surfaces in the Al-5% Mg-0.1 ppm Na alloy fractured by liquid gallium (Fig. 5(a)), whereas several dimples were observed on the GB surface
in the Al-5% Mg-2 ppm Na alloy (Fig. 5(b)). Below 673 K, the fraction of GB sliding in the impact test was sub-stantially smaller than that in the quasi-static test. Since GB sliding is caused by lattice diffusion around the GBs, it is assumed that such lattice diffusion would not occur in the impact test. This is because the time required for the impact fracture was very short (about 300ms at 573 K) in comparison with the quasi-static fracture (about 90 s at 573 K).
In addition, it was observed that the morphology of the serrated GBs in the Al-Mg alloys deformed at elevated temperatures. This is indicative of an initial state of GB
Fig. 6 SEM micrographs of the serrated grain boundaries tested at elevated
temperatures. Figures 6(a) and (c) show the micrographs for the Al-5% Mg-0.1 ppm Na alloy tested by the quasi-static test and the impact test, respectively; (b) and (d) show the micrographs for the Al-5% Mg-2 ppm Na alloy tested by the quasi-static test and the impact test, respectively.
Fig. 7 Schematic diagrams of high temperature embrittlement in the
[image:4.595.202.519.67.572.2] [image:4.595.56.300.71.603.2]GB migration barely occurred at elevated temperatures in the impact test. On the basis of such microscopic differences, it is believed that the improvement of the ductility in the quasi-static deformation above 673 K was caused by dynamic recovery.
4. Summary
The HTE phenomenon in the Al-5% Mg alloy is summa-rized as shown Fig. 7. In the quasi-static deformation, HTE is caused by GB fracture, which is attributed to the nucleation, growth, and coalescence of cavities at the GBs. The stress concentration at the GBs is caused by GB sliding accompanied by GB serration around 573 K. It is assumed that the segregation of sodium at the GBs promotes GB sliding during the quasi-static test. It is believed that the improvement in the ductility above 673 K is caused by GB migration; the improvement is attributed to dynamic recovery, which reduces the stress concentration at the GBs. In the impact deformation, on the other hand, HTE is caused by GB fracture. The impact deformation is attributed to the slip localization at the GBs weakened by the segregation of sodium. The low fraction of GB sliding and the GB migration in the impact test would prevent dynamic recovery at 573 K and 673 K, where the effect of the segregation of sodium at the GBs on changes on the microscopic scale was small.
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