0 Introduction: Some Economic Aspects of Nonferrous Metals Karl Heinz Matucha
Metallgesellschaft AG, Frankfurt am Main, Germany
List of Symbols 0.1 0.2 0.3 0.3.1 0.3.2 0.3.3 0.3.4 0.4 0.5 0.6 0.7
Smelter Production Consumption and Scrap Recovery Detailed Consideration of Selected Metals Aluminum Copper Zinc Lead Reasons for the Variation of Production and Consumption Scrap Recovery and Recycling The Future of Nonferrous Metals References
Materials Science and Technology Copyright © WILEY-VCH Verlag GmbH & Co KGaA. All rights reserved.
2 2 3 4 5 8 13 15 16 19 21 22
0 Introduction: Some Economic Aspects of Nonferrous Metals
List of Symbols
the development and production of highalloy steels and superalloys, the metals with a low density (aluminum, magnesium, titanium) obtained their increasing importance on the basis of their own specific properties. Since 1960 aluminum has become the most important nonferrous metal, followed by copper, zinc and lead. In addition to technological reasons, the worldwide production, manufacture and consumption of metals depends also on the general economic situation. Therefore the smelter production of nearly all metals decreased in the period from 1928 to 1933. On the other hand, the end of world war II also resulted in a decrease in the smelter production, especially of iron, aluminum and magnesium. This breakdown of production was followed by a high yearly increase in production. Independently of the fluctuating general economic situation, the median yearly
utilization potential recycling rate
0.1 Smelter Production The smelter production of iron and some nonferrous metals between 1880 and 1980 in general showed an increase in production for all materials (Fig. 0-1). It should be noted, however, that the production of iron was much larger than the total production of all nonferrous metals in Fig. 0-1. Of the nonferrous metals, nickel, aluminum, magnesium and titanium had the highest growth rates in production (Matucha and Wincierz, 1986). While the growing need for nickel has been a result of
Fe (raw steel)
10V
108
to7
106
105 o
104
Al/
1880
90 1900
Figure 0-1. Smelter production (total world) of important metals from 1880 to 1980 (Matucha and Wincierz, 1986).
Mg/
10
20
30 Year
40
50
60
70
1980
0.2 Consumption and Scrap Recovery
rates of growth of smelter production have shown a falling tendency for the last thirty years. The yearly growth rates in the period 1971-1980 were smaller than in the previous decade (Matucha and Wincierz, 1986). This tendency continued during the decade 1981-1990 (Table 0-1). For aluminum, copper, lead, zinc, tin and nickel, the yearly growth rates decreased considerably, leading to "negative" growth for lead and tin. The numbers given in Table 0-1 are world average values. For a detailed consideration it would be necessary to look at the actual yearly smelter production (Metallgesellschaft, 1989, 1994). This has
Table 0-1. Yearly growth rates (%) of consumption, smelter production and scrap recovery (Western world). 19711981
19811991
+ 13.2 Al: Consumption Smelter Production + 13.9 + 13.3 Scrap Recovery
+ 4.7 + 4.5 + 8.8
+ 2.8 + 1.7 + 4.5
Cu: Consumption Smelter Production Scrap Recovery
+ 4.1 + 4.4 + 4.2
+ 3.6 + 2.8 + 1.8
+ 1.0 + 0.9 + 2.5
Pb: Consumption Smelter Production Scrap Recovery
+ 4.4 + 4.1 + 2.7
+ 3.5 + 0.5 + 6.3
+0.3 -0.7 + 2.1
Zn: Consumption Smelter Production Scrap Recovery
+ 5.5 + 4.9 + 2.5
+ 2.3 + 2.6 + 5.9
+ 1.2 + 1.5 + 0.9
Ni: Consumption Smelter Production Scrap Recovery
+ 6.1 + 7.0 n.k.a
+ 2.6 + 1.6 n.k.a
+ 2.4 + 2.0 n.k.a
Mg: Consumption + 8.3 Smelter Production + 11.8 Scrap Recovery + 7.7
+ 1.5 + 3.1 + 22.2
+ 1.9 + 1.6 + 0.07
-1.3 + 0.3 -1.0
+ 0.5 -0.9 -0.8
19611971
Sn: Consumption Smelter Production Scrap Recovery a
n.k.: not known.
+ 2.0 + 2.0 -1.4
—a
t i
Al
Pb
106 a —a—
a
—a—a—a— a
Ni
3 o O—u—•—D—•—•—o—•—•—D
105 1978 79 80 81 82 83 84 85 86 87 88 89 90 91 92 1993 Year
Figure 0-2. Smelter production (total world) of some important nonferrous metals from 1978 to 1993.
been done for the period from 1978 to 1993 for aluminum, copper, zinc, lead, nickel, tin and magnesium (Fig. 0-2). The poor economic situation at the beginning of the 1980s led to a decrease in smelter production of aluminum, copper, zinc, nickel and magnesium, whereas the smelter production of tin and lead was nearly constant. After 1982 the smelter production generally increased. There has been a smaller increase for aluminum, copper and tin since 1989, while the production of nickel and magnesium has decreased slightly.
0.2 Consumption and Scrap Recovery As can bee seen in Table 0-1, the total consumption of the nonferrous metals had the highest yearly growth rates between
0 Introduction: Some Economic Aspects of Nonferrous Metals
smelter production. This difference was highest for lead, followed by copper and aluminum. The scrap recovery (Fig. 0-4) indicates high growth rates for aluminum since 1984. The reasons for this and other influences on scrap recovery will be discussed below.
1961 and 1971. These growth rates in general decreased during the following decade, showing considerable differences between the metals. For the scrap recovery of the metals data from Western countries only are available. The data compiled in Table 0-1 show that the scrap recovery of the metals - with the exception of Sn - has grown, too. It should be noted that during the last decade (1981-1990) the yearly growth rates of scrap recovery for aluminum, copper and lead exceeded those of the smelter production. The total consumption of aluminum, copper, lead, nickel and magnesium since 1978 is given in Fig. 0-3, together with the consumption of primary zinc and tin. Comparison with Fig. 0-2 shows that the development of the smelter production followed that of the consumption. In nearly all cases the consumption exceeded the
0 3 Detailed Consideration of Selected Metals The global data on the primary production, consumption and secondary production discussed so far could lead to the conclusion that during the last few years no remarkable changes have occurred. To prove that conclusion, the primary production, consumption and secondary production will be analyzed in detail for the four most important nonferrous metals
Cu _m-—in—ts—B—CD—m—u
Zn
_E—B—B—B—B—B—B'—E—B—E Ph
.9
10 6
Q.
-H—a—a" O
o
10 5
1978 79 80 81 82 83 84 85 86 87 88 89 90 91 92 1993 Year
Figure 0-3. Total consumption of important nonferrous metals from 1978 to 1993.
0.3 Detailed Consideration of Selected Metals ,a—n Al
si—Bl—IB—SI—OB—™—«"—™
Zl1
Figure 0-4. Scrap recovery of aluminum, copper, lead and zinc since 1978. Western countries (Metallgesellschaft, 1989, 1994). 1978 79 80 81 82 83 84 85 86 87 88 89 90 91 92 1993 Year
(aluminum, copper, zinc and lead) since 1978. To do this, the worldwide data are divided according to geographical region. Furthermore, the so-called "Western" and "Eastern" countries are treated separately. (In the Metallstatistik (Metallgesellschaft, 1989, 1994) all communist and formerly communist states are included under "Eastern countries", regardless of their geographical location.) Additionally, the main producer and consumer are considered.
the beginning of the 1980s and in 1985 to 1987. On the other hand, production in America as a whole has increased sharply since 1986. This results from the increasing production of primary aluminum in Canada, Brazil and Venezuela. Canada has become number three in the world (Table 0-2) while Brazil and Venezuela entered the list of the top ten and improved their positions during the last few years.
0.3.1 Aluminum
Table 0-2. Primary aluminum: Main producer countries and their contributions (%) to the total production in 1978 and 1993.
Primary Production
1978
(%)
1993
(%)
The main producers of primary aluminum are given in Table 0-2 for 1978 and 1993, together with the total production. More than 59% of the total production was produced by five countries, while more than 72% of the total production came from ten countries. The production of these ten states is listed in Table 0-2 for each year. The United States remained the largest producer of primary aluminum, although the production for some years since 1978 shows a slightly falling tendency (Fig. 0-5) driven by the poor economic situation at
U.SA U.S.S.R. Japan Canada Germany (West) Norway France P.R. China U.K. Australia
28.2 19.4 6.8 6.8 4.8 4.2 2.5 2.3 2.2 1.7
U.S.A. C.I.S. Canada Australia P.R. China Brazil Norway Venezuela Germany India
18.9 15.6 11.8 7.0 6.2 6.0 4.5 2.9 2.8 2.4
Total world (100%): (103 tonnes)
78.9
78.1
14767
19600
0 Introduction: Some Economic Aspects of Nonferrous Metals 8000 r ^ A m e r i c a (incl. U.S.A.)
7000 -
Eastern states
e Europe
New Zealand and Australia ^ " (excl. Eastern states) ,.—©—-©—©—©—©—©~""'P,. Africa
Figure 0-5. Primary production of aluminum since 1978 in different regions of the world. Production in the United States is plotted separately (Metallgesellschaft 1989, 1994).
1978 79 80 81 82 83 84 85 86 87 88 89 90 91 92 1993 Year
The production of primary aluminum in the Eastern bloc, which is partly estimated, showed small variations and did not depend on the general economic situation. The increase of the production since 1983 is mainly due to the increasing production in the People's Republic of China. In Europe (excluding the Eastern states), production was nearly constant from 1978 to 1991 and decreased afterwards. Taking into account that the total production has increased it follows that the proportion of the European primary aluminum contribution to total world production decreased from 25% in 1978 to 17% in 1993. This follows qualitatively from Table 0-2, too, which shows that there were only two European producers among the top ten in 1993. In Australia and New Zealand the production of primary aluminum has grown since 1983; Australia's production has increased since 1983 by 19% per year.
All producers of primary aluminum in Africa (Egypt, Ghana, Cameroon, South Africa) have increased production, albeit at a low level. There the growth rate of Ghana is the highest. Production in Asia (excluding the Eastern states) showed a decrease from 1978 to 1983 and then an increase. The reasons for this behavior follow immediately from Table 0-3. In 1978, 70% of Asia's primary
Table 0-3. Production of primary aluminum (in thousand tonnes) in Asia (excluding Eastern states). 1978
1983
1988
1993
Japan India Bahrain Indonesia United Arab Emirates
1057.7 205.4 122.8 _
255.9 204.8 171.7 114.8 151.2
35.3 334.5 182.8 185.1 162.5
18.3 466.4 450.0 202.1 242.3
Total Asia
1511.9
980.6
1013.0
1529.1
0.3 Detailed Consideration of Selected Metals
aluminum was produced in Japan. Production in Japan was then nearly stopped, whereas India and Bahrain increased production. Indonesia and the United Arab Emirates became producers of primary aluminum. Consumption of Primary Aluminum In comparison to the production of primary aluminum the consumption shows smaller variations in the list of the top ten consumers in 1978 and 1993 (Table 0-4). The changes which occurred between 1988 and 1993 are more clearly visible in the yearly consumption of primary aluminum (Fig. 0-6). This yields two main features: - a drastic decrease in consumption in the Eastern states although the consumption in P.R. China increased; - an increase in consumption in Asia overall. A detailed consideration of the consumption of primary aluminum in Asia (Table 0-5) shows that the proportion of
Table 0-4 Primary aluminum: Main consumer countries and their proportions (%) of the total consumption in 1978 and 1993. 1978
(%)
1993
(%)
U.S.A. U.S.S.R. Japan Germany (West) France P.R. China Italy U.K. Canada Belgium/ Luxembourg
32.4 11.9 10.8 6.2 3.4 3.3 2.6 2.6 2.2 1.7
U.S.A. Japan P.R. China CIS. Germany France Korea Italy Canada U.K.
26.3 11.7 7.1 6.4 6.3 3.6 3.0 3.0 2.7 2.6
Total world (100%): (103 tonnes)
77.1
72.7
15 348
18 541
Japan's consumption decreased. Other countries increased their consumption considerably. The percentage of the consumption of the main consumer in Asia (Japan) decreased from 71% in 1978 to
7000 r America 6000 L \
,o—-o—o'
-Europe
5000 \ 4000
/ *
—e—e—e—*
Asiaflncl. Japan)
Eastern s t a t e s \ (incl. P.R. China) ®
3000
g 2000\o O
Q
~~Q
P.R. China •®—®~~®
Australia Africa
1978 79 80 81 82 83 84 85 86 87 88 89 90 91 92 1993 Year
Figure 0-6. Consumption of primary aluminum since 1978 in different regions of the world. The consumptions in Japan and P.R. China are plotted separately (Metallgesellschaft 1989, 1994).
8
0 Introduction: Some Economic Aspects of Nonferrous Metals
Table 0-5. Consumption of primary aluminum (in thousand tonnes) in Asia (excluding Eastern states). 1978
1983
1988
1993
Japan India Korea Taiwan Iran Turkey Thailand Bahrain
1656.1 205.4 105.8 89.9 53.4 45.0 33.7 7.1
1722.0 218.5 120.0 136.5 105.0 89.2 64.9 21.3
2123.2 327.0 268.0 175.7 100.7 109.7 71.7 81.0
2174.8 475.3 557.9 299.1 110a 128.6a 177.4 124.7 a
Total Asia
2327.6
2682.8
3533.4
4554.6
1992.
48% in 1993, indicating an increase in consumption in other Asian countries. The gap in the consumption between Europe and Asia was (Fig. 0-6) nearly constant between 1978 and 1986. It became smaller after 1986, and in 1993 the consumption of primary aluminum in Asia and Europe was nearly equal. In America as a whole large variations of the consumption are visible. It should be noted that the consumption in 1993 was nearly the same as in 1979, whereas in the intervening period the consumption was less. In Africa and Australia only small amounts of aluminum were consumed. Scrap Recovery of Aluminum (Secondary Aluminum) Only data for the Western countries are available (Fig. 0-7). They show, in comparison to Fig. 0-4, that the total increase of the production of secondary aluminum is very largely due entirely to the United States. It should further be noted that in Europe since 1989 and in Japan since 1991 the production of secondary material has decreased. Japan is the major producer of secondary aluminum in Asia but during the last few years Taiwan, which has
stopped production of primary aluminum has increased secondary production. Total Consumption of Aluminum Comparison of the total consumption of aluminum in the world (Fig. 0-3) with the consumption in the different regions (Fig. 0-8) shows that variations of the total consumption are mainly determined by the consumption of the Western countries in Asia, America and Europe. Consumption in Africa and Australia has been nearly constant since 1978, while consumption in the Eastern states, having increased slightly until 1990, dropped in 1991. Consumption in the Western world increased after the period from 1979 to 1982. In the following years (1983-1989) consumption showed a steady growth, especially in Europa, whereas large growing rates are observed in Asia from 1986 to 1991. This led to a smaller difference in consumption between Europe and Asia. Although Japan has remained the main producer in Asia, it can be seen that other Asian countries increased their consumption of aluminum. In Korea consumption increased by a factor of eight from 1980 to 1993. As a consequence Korea has become one of the main consumers of aluminum. 0.3.2 Copper Smelter Production from Ores According to Fig. 0-2 the total smelter production of copper has shown a nearly continuous increase since 1978. Nevertheless, there have been exceptional changes in production in different regions of the world (Fig. 0-9). These occurred from approximately 1986 and can be characterized by a decreasing production in Africa and growing productions in Asia and America, while production in the Eastern countries
0.3 Detailed Consideration of Selected Metals America (incl. U.S.A.K® ?
3000
U.S.A.
2500
-ST2000 0
e
y
D—(D
Asia
0 ©—Q
8 1000
Japan
Figure 0-7. Scrap recovery of aluminum (secondary aluminum) in the Western countries since 1978. Scrap recovery in the United States is plotted separately (Metallgesellschaft 1989, 1994).
o
CO
500f
1978 79 80 81 82 83 84 85 86 87 88 89 90 91 92 1993 Year
-—7
9000
U.S.A.
8000
7000
g
6000
6
5000
"«»
y
/
s—
«'
V s-~
Europe —(D—(D
Asia (incl. Japan)
Eastern states
Japan
3000—©—©—9' 2000
1000 /
Australia Africa
1978 79 80 81 82 83 84 85 86 87 88 89 90 91 92 1993 Year
Figure 0-8. Total consumption of aluminum since 1978 in the different regions of the world. The consumptions in the United States and Japan are plotted separately (Metallgesellschaft 1989, 1994).
10
0 Introduction: Some Economic Aspects of Nonferrous Metals
4000
500G—e
I
^Australia
Figure 0-9. Smelter production of copper from ores since 1978 in the different regions of the world. The productions in the United States and in Japan are plotted separately (Metallgesellschaft 1989, 1994).
1978 79 80 81 82 83 84 85 86 87 88 89 90 91 92 1993 Year
decreased from 1988 to 1991 and increased after 1991 due to the fact that more copper was produced in P.R. China and Poland (Table 0-6). In Africa all copper-producing countries reduced production after 1986/1987. This reduction was pronounced for Zaire, where in 1993 only 10% of the 1983 production - when Zaire was number six in the list of the top ten producers - was achieved. The production of copper from ores was nearly the same in 1978 and 1993 in the United States; the growing production in America results from the growing production in the other countries (see Table 0-7). In Asia, Japan is the largest producer of copper from ores, but other countries, e.g. Iran, Oman, Korea and the Philippines,
Table 0-6. Copper: Smelter production from ores. Main producer countries and their contributions (%) to the total production in 1978 and 1993. 1978
(%)
1993
(%)
U.S.A. U.S.S.R. Chile Japan Zambia Canada Zaire Poland Peru South Africa
16.7 14.6 12.0 11.0 8.5 5.4 5.1 4.1 4.1 2.7
Chile U.S.A. Japan CI.S. Canada Zambia P.R. China Poland Australia Peru
15.1 13.7 11.8 9.8 5.8 4.9 4.8 4.5 3.5 3.4
Total world (100%): (103 tonnes)
84.2
77.3
7735
9213
11
0.3 Detailed Consideration of Selected Metals
Table 0-7. Copper: Smelter production from ores (in thousand tonnes) in America. 1978
1983
1988
1993
926.6 1288.4 420.6 319.0 85.9 -
1058.9 927.7 405.7 295.9 59.5 63.1
1189.4 1043.0 478.3 239.2 150.3 147.9
1389.1 1265.2 535.9 317.2 281.5 150.0
Total America 3041.0
2811.0
3248.0
3939.0
Chile U.S.A. Canada Peru Mexico Brazil
have increased their production during the
Production of Refined Copper The production of refined copper exceeded the production of copper from ores; the difference between these productions has become larger since 1978 according to Tables 0-6 and 0-8. With respect to the different regions, production of refined copper (Fig. 0-10) shows the same variations as production from ores. It should be noted, however, that in Europe and in America the production of refined copper is high in comparison to the production of copper from ores. In America this results mainly from the increased production of refined copper in the United States and Chile (see Table 0-9).
Table 0-8. Copper: Refined production. Main producer countries and their contribution (%) to the total production in 1978 and 1993. 1978
(%)
1993
U.S.A. U.S.S.R. Japan Chile Zambia Canada Germany Poland P.R. China Zaire
20.3 14.2 10.6 8.3 7.0 4.9 4.5 3.7 2.9 2.7
U.S.A. Chile Japan C.I.S. P.R. China Germany Canada Zambia Poland Belgium/ Luxembourg
Total world (100%): (103 tonnes)
20.3 11.4 10.7 7.2 6.2 5.7 5.1 3.8 3.6 3.0
79.1
77.0
9030
11092
Table 0-9. Copper: Refined production (in thousand tonnes) in America. 1978
1983
1988
1993
Chile U.S.A. Canada Peru Mexico Brazil
748.2 1832.0 446.3 185.6 83.0 25.9
834.2 1583.8 464.3 194 A 83.5 63.1
1012.7 1852.4 528.7 179.6 140.8 147.9
1268.2 2252.5 561.6 261.7 181.0 158.4
Total America
3331.0
3235.0
3878.0
4702.0
Scrap Recovery The data for scrap recovery in the Western countries (Table 0-10) summarize the amounts of refined copper obtained from scrap and the scrap used directly in manufacture. The preparation of refined copper makes up nearly 30% of the total scrap recovery. Approximately 70% of the scrap was recovered by just five countries and more than 80% by the ten countries in Table 0-10. The total scrap recovery of
copper increased considerably between 1978 and 1988. Taking into account the total scrap recovery data, it follows from Table 0-10 that scrap recovery has increased considerably in Germany, Italy and Japan. Since 1983, Korea has been one of the leading countries for scrap recovery of copper.
0 Introduction: Some Economic Aspects of Nonferrous Metals
12
4500
4000
(incl. U.S.A.)
3500
3000
§2500 o
\
f-k—r±—/k.—<|—
Eastern states *<
U.S.A. $
c 2000 h, o c
©
•D
Asia Europe
2 1500F ~e—©—c ° \
_^ Africa Australia
500 I I
i
i
Figure 0-10. Production of refined copper since 1978 in the different regions of the world. The production in the United States is plotted separately (Metallgesellschaft 1989, 1994).
1978 79 80 81 82 83 84 85 86 87 88 89 90 91 92 1993
Year
Table 0-10. Copper: Scrap recovery (Western world). Main countries with their contribution (%) to the total scrap recovery. 1978
(%)
1993
U.S.A. Japan Germany (West) U.K. Italy France Yugoslavia Belgium/ Luxembourg Australia Canada
38.5 14.2 8.1 6.9 6.2 4.1 2.6 2.6 1.7 1.6
U.S.A. 24.8 Japan 15.7 Germany 14.9 Italy 7.6 Korea 3.9 France 2.9 U.K. 2.4 Belgium/ 2.2 Luxembourg Australia 1.6 Spain 1.5
86.5
81.1
3299
4697
Total world (100%): (103 tonnes)
(%)
Consumption of Refined Copper Since 1978 the total consumption of refined copper has grown (Fig. 0-3) with small variations. In the different regions of the world the variations in consumption have been very large, however (Fig. 0-11). The main feature is the increasing consumption in Asia. This resulted mainly from the growing consumption in Korea and Taiwan (Table 0-11), but in other Asian countries, e.g. Thailand, Saudi Arabia, Malaysia, Indonesia, Iran, Turkey and India, consumption of refined copper has increased, too. In comparison to Asia, consumption in Europe has shown smaller growth rates, whereas consumption in America is characterized by large variations. While in 1978 consumption in Asia was only ~ 5 0 % of the consumption in
0.3 Detailed Consideration of Selected Metals
13
©_-©—©—©—.©^ yXsiaOnd
ft
Figure 0-11. Consumption of refined copper in the different regions of the world since 1978. The consumption in Japan is plotted separately (Metallgesellschaft 1989, 1994). 1978 79 80 81 82 83 84 85 86 87 88 89 90 91 92 1993 Year
Table 0-11. Refined copper: Main consumer countries and their proportions (%) of the total consumption in 1978 and 1993. 1978 U.S.A. U.S.S.R. Japan Germany (West) U.K. P.R. China Italy France Belgium/ Luxembourg Canada
Total world (103 tonnes)
1993 23.0 13.9 13.0 8.2 5.3 3.7 3.6 3.3 3.0
U.S.A. Japan P.R. China Germany CIS. Italy Taiwan France Korea U.K.
21.6 12.6 8.6 8.1 5.0 4.4 4.3 4.3 3.6 3.0
2.6 79.6
75.5
9 535
10980
Europe or in America, in 1993 the consumption of refined copper was nearly the same in America, Europe and Asia. The
large decrease in consumption in the Eastern states has been partly balanced by the growing consumption in P.R. China (Table 0-11). 0.33 Zinc Smelter Production The slow but continuous increase of the total smelter production (Fig. 0-2) since 1982 results from the different production data in the world (Fig. 0-12). The main feature is the high growth rate of the production in P.R. China, which has become the main producer of zinc, followed by Japan and Canada. Most of the smelter production of zinc is done in Europe, with Germany, Spain, France and Italy as the main producers. Although the main producer in Asia, Japan, reduced production, the total production in Asia has increased since 1982 because Korea, India and Thailand enhanced their production.
14
0 Introduction: Some Economic Aspects of Nonferrous Metals
2000-
1500Eastern states (incl. P.R. China) CD C
I .1000
500 -
Figure 0-12. Smelter production of zinc since 1978 in the different regions of the world. The production in Japan and P.R. China are plotted separately (Metallgesellschaft 1989, 1994). 1978 79 80 81 82 83 84 85 86 87 88 89 90 91 92 1993
Year
Consumption of Primary Zinc
Scrap recovery
The consumption of primary zinc in America has shown large variations since 1978 (Fig. 0-13), being affected mainly by the values of the leading consumer in this region (U.S.A.). While in Europe consumption increased from 1982 to 1991, it was nearly constant in Japan during this time. However, consumption in Asia grew during this period because other countries, e.g., Korea, Taiwan and Thailand, consumed more primary zinc. Consumption in the Eastern countries decreased considerably.
The total scrap recovery of zinc (Fig. 0-4) can be divided into the following groups: -
scrap recovery in primary zinc smelters remelted zinc and remelted zinc alloys zinc in copper and other alloys direct use of scrap in chemicals, paints, pigments, etc.
The data available for the four groups are incomplete, but the following relationships can be estimated for 1983-1993. The scrap recovery in primary zinc smelters was nearly constant, with 30% of the total scrap recovery, whereas 15-20% of the
0.3 Detailed Consideration of Selected Metals
15
2000
-Z
\
Asia (incl. Japan)
/
e
1500 \
/ America
o—q \
CD C O
Eastern states (incl. P.R. China)
1000
^ N ^
Japan
500
P.R. China
Africa
Australia
-I 1 I L_ _J L _J L. J L 1978 79 80 81 82 83 84 85 86 87 88 89 90 91 92 1993 Year
total scrap recovery was by remelting operations. The direct use of scrap in chemicals etc. was in the range of 20%. The remainder was used in copper and other alloys. 0.3.4 Lead The smelter production of lead in the different regions of the world (Fig. 0-14) has been marked by pronounced changes: a decrease in production in America and an increase of production in Asia. On the other hand, the production of refined lead (including secondary material) has increased in America since 1983 (Fig. 0-15). The same is true for production in Asia, where other countries beside Japan increased production of refined lead, e.g., Korea, Taiwan, and India.
Figure 0-13. Consumption of primary zinc since 1978 in the different regions of the world. The consumption in Japan and P.R. China are plotted separately (Metallgesellschaft 1989, 1994).
The total consumption of lead increased from 1983 to 1989 (Fig. 0-3). This period was followed by years with slightly decreasing consumption. Figure 0-16 shows that the growing consumption (1985 — 1989) was caused by more consumption of lead in America, Europe, and Asia. The enhanced consumption in Asia since 1982 indicates that Western Asian countries (except Japan) have increased their consumption of lead. This can be deduced from Table 0-12. The proportion of refined lead obtained by scrap recovery in the Western countries has been high, at 50% in 1978, 44.3% in 1983 and 53.3% in 1988 and 1993.
16
0 Introduction: Some Economic Aspects of Nonferrous Metals
'[A
America (incl. U.S.A.)
urope
Africa
Figure 0-14. Smelter production of lead since 1978 in the different regions of the world. The production in the United States and in Japan are plotted separately (Metallgesellschaft 1989, 1994).
1978 79 80 81 82 83 84 85 86 87 88 89 90 91 92 1993 Year
Table 0-12. Consumption of refined lead (in thousand tonnes) in Asia (excluding Eastern states). 1978
1983
1988
1993
Japan Korea Taiwan India Thailand Indonesia Malaysia
352.1 27.4 24.2 55.0 8.1 8.6 2.8
539.6 57.5 39.5 60.0 22.3 11.8 11.3
406.5 146.0 74.8 75.0 28.5 17.0 23.9
371.3 177.0 100.7 80.0 50.0 75.0 65.0
Total Asia
535.7
654.7
880.3
1054.4
0.4 Reasons for the Variation of Production and Consumption The analysis of the production and consumption data for the four nonferrous metals shows differences in the total general data and more pronounced variations in the data of the different regions of the world. With respect to the global data the high but strongly changing growth rate in the consumption of aluminum and the slightly decreasing but nearly constant use of lead are the main topics. These differences result from changes in the fields of applications. This can be seen in Figs. 0-17 and 0-18, which show the
0.4 Reasons for the Variation of Production and Consumption
17
America (incl. U.S.A.)
2000 r
\
/
v
e
1500
\
Europe
e—e
"V J
.S.A.
Eastern states
g
1 I a. Asia (incl. Japan)
500
Japan
Q—9—©—
Australia #
©-
Africa _l
I
I
Figure 0-15. Production of refined lead since 1978 in the different regions of the world. The production in the United States and in Japan are plotted separately (Metallgesellschaft 1989, 1994).
1978 79 80 81 82 83 84 85 86 87 88 89 90 91 92 1993 Year
main fields of application of aluminum for the United States and Japan. Owing to steadily increasing use, in the United States packaging has become the main field of application, which seems to be independent of the economic situation, in contrast to the applications in transportation and building and construction. In Japan, transport has become the main field of application, but the use for packaging has also increased. These tendencies are in general also found in other countries. Especially for building, construction and packaging the growth rates have been high. The situation for the applications of lead is different. While the use of of lead in
the main field of application (batteries) has increased, other fields of application, e.g. cables and in particular tetraethyl lead (antiknock) compounds, have lost their importance. While in 1978 nearly 50% of lead was used for batteries in the U.S.A., Japan, Germany, Italy, France and United Kingdom, the application of lead for batteries was 66% of the consumption in 1992. Several reasons are responsible for the variations of production and consumption of all metals in the different regions of the world. Political reasons: These are very clearly visible in the development of the data for
18
0 Introduction: Some Economic Aspects of Nonferrous Metals 2000 r America (incl. U.S.A.)
I
I
1978 79 80 81 82 83 84 85 86 87 88 89 90 91 92 1993 Year
Figure 0-16. Consumption of refined lead since 1978 in the different regions of the world. The consumptions in Japan and the United States are plotted separately (Metallgesellschaft 1989, 1994).
Packaging Transport •
1000
Building and construction Electrical engineering
v
Household articles
1978 79 80 81 82 83 84 85 86 87 88 89 90 91 92 1993 Year
Figure 0-17. Consumption of aluminum in the United States in main fields of application since 1978 (Metallgesellschaft 1989, 1994).
0.5 Scrap Recovery and Recycling
/
1000
19
Transport and construction
^
500 Packaging
3
8
Electrical engineering
Figure 0-18. Consumption of aluminum in Japan in main fields of application since 1978 (Metallgesellschaft 1989, 1994).
1978 79 80 81 82 83 84 85 86 87 88 89 90 91 92 1993 Year
the Eastern countries, which showed nearly constant values until 1989/90. The decrease of the figures thereafter is pronounced for consumption and less pronounced for production. The resulting higher rates of export led to a breakdown of the metal prices and to difficulties for the producers in the Western world. Economic reasons: Because the consumption of metals depends on the degree of industrialization, the increase of the use of metals will have small growth rates in the industrial states if no new fields of application exist (see the special case of consumption of aluminum, however). On the other hand, countries developing into industrial countries have high growth rates in the consumption of metals. In a second stage these countries will start or increase the production of metals, too, if other conditions, e.g., availability of energy and raw materials, are fulfilled. Typical examples of this development have been given above. The primary production of metals from raw materials is an energy-consuming process, energy consumption being extremely high for aluminum. Therefore countries
with high energy prices reduced primary production and increased secondary production. Japan is a typical example. There, primary production was nearly stopped while secondary production was increased, with the result that ~ 30% of Japan's aluminum requirement is produced in secondary smelters. On the other hand, countries with a high output of energy, e.g., Venezuela and the United Arab Emirates, became important producers of primary aluminum. In this context it should be noted that in Europe Norway was the only country with high growth rates in the production of primary aluminum. The trend to produce more secondary material will continue in the highly industrialized countries, which are the main producers and in some cases have no or only small resources of raw materials. Therefore the old scrap will become an important source for production.
0.5 Scrap Recovery and Recycling Since the report of the Club of Rome ("Limits of Growth") appeared in 1973,
20
0 Introduction: Some Economic Aspects of Nonferrous Metals
'B—E—B Pb
Figure 0-19. Proportion of scrap and residues in the total consumption in the Western world. For aluminum the data of the United States are plotted separately (Metallgesellschaft 1989, 1994). 1978 79 80 81 82 83 84 85 86 87 88 89 90 91 92 1993 Year
the discussion related to the recovery of materials and the use of primary energy has gained in significance. In contrast to these discussions, the growth rates of the production of recovered materials - except aluminum - seem to be very small (Fig. 0-7). This impression is strengthened when scrap and residues are considered as a proportion of the total consumption of some nonferrous metals in the Western world (Fig. 0-19). The proportions are quite different for the various metals, with high amounts for lead followed by copper, aluminum, and zinc. Furthermore, the proportions vary as a function of year. The development since 1978 shows a slightly increasing tendency for aluminum and for lead since 1983. Two main reasons determine the use of recovered materials as a fraction of the total consumption. 1. In the Western world the recovery of materials was carried out mainly under economic aspects. If the price of the primary metal was high, the proportion of the recovered metal was high, too, and vice versa. Therefore, in periods of increasing prices of the primary metals,
the proportion of scrap recovery increased. This was, for example, the case for copper from 1978 to 1980 and for lead from 1986 to 1990. Therefore the proportion of scrap recovery in the total consumption varied from year to year. 2. The amount of scrap recovery depends on the fields of application: Because copper is mainly used as the pure metal and as alloys with high copper content, the recovery is easy. On the other hand, zinc is mostly used in galvanizing, which leads to difficulties with respect to recovery. A further fraction of zinc is used for chemicals and cannot be recovered. The high proportion for lead and the increasing tendency result from the increasing use of lead for batteries, where the lead can be recovered. The scrap recovery of aluminum (Fig. 0-9) is high in the United States, and the proportion of scrap-recovered aluminum in the U.S.A. is higher than in other Western countries (Fig. 0-19). This results from a special development: In the United States the use of aluminum for packing (cans) has increased since the middle of the 1970s (see also Fig. 0-17). Furthermore, a system for collecting used alu-
0.6 The Future of Nonferrous Metals
minum cans has been introduced, with the result that the recovery rate of the cans has increased. In principle, the potential for recovery is based on new scrap, residues, and old scrap. While in general the volume of new scrap is proportional to the manufactured volume of a metal, the amount of residues will increase because emissions must be reduced. Due to the increasing costs for the deposition of residues, more residues will be recovered. The potential volume of old scrap results from the consumption of metals five, ten or twenty years previously, because utility goods have lifetimes of five years or more, e.g., automobiles 10 years, cables 30-40 years. Taking into account the higher growth of the consumption in the past (Table 0-1), the potential of old scrap will increase for statistical reasons alone. This will happen in those countries which are the main consumers of the metals. Some of these countries have no or only few resources of raw materials. Therefore the only source for the production of metals will be scrap.
0.6 The Future of Nonferrous Metals Independently of the economic situation, the analysis of the past allows some conclusions to be drawn for the future. - The nonferrous metals will retain their main fields of applications because of their specific properties. - New fields of application will arise because the consumption of energy and emissions from machinery must be reduced. There will be a need, for example, for improved high-temperature materi-
21
als and improved materials with low densities. - The total consumption of nonferrous metals will increase because the degree of industrialization in developing countries will increase. - In the competition between different metals and other materials, ecological aspects, including the possibility of recycling, will play an important role. - An important advantage of metals is the possibility of recycling. For this, appropriate technologies must be available together with collection systems. Finally, it should be pointed out that even small increases in the amounts of recovery and recycling will be important for the use of the materials. Therefore, we define the potential of utilization to be equal to one for a single use (recycling rate = 0) and infinite for a recycling rate of 1, corresponding to 100% recycling. The utilization potential Pp is then given by 1 (0-1) 1-x where x is the recycling rate ( 0 < x < l ) . If a recycling rate x0 is increased to x0 + Ax the corresponding increase of the utilization potential is 1
Ax
(0-2)
From Eqs. (0-1) and (0-2) the following conclusions can be drawn: - For small recycling rates the utilization potential is small. - The utilization potential increases more than linearly with an increasing recycling rate. - The higher the recycling rate x the higher the increase of the utilization potential created by small increases in recycling rates.
22
0 Introduction: Some Economic Aspects of Nonferrous Metals
0.7 References
General Reading
Matucha, K. H., Wincierz, P. (1986), in Chemische Technologic Vol. 4, 4th ed.: Hamisch, H., Steiner, R., Winnacker, K. (Eds.) Munich: Carl Hanser, pp. 1-37. Metallgesellschaft (Ed.). (1989), Metallstatistik 1978-1988, 76th ed. Frankfurt: Metallgesellschaft. Metallgesellschaft (Ed.) (1994), Metallstatistik 19831993. 81st ed. Frankfurt: Metallgesellschaft and World Bureau of Metal Statistics.
Metallgesellschaft (Ed.), Metallstatistik. Frankfurt: Metallgesellschaft. A volume has appeared each year since 1978.
1 Tin Ashok Kumar Suri and Srikumar Banerjee Metallurgy Division, Bhabha Atomic Research Centre, Trombay, Bombay, India
List of 1.1 1.2 1.2.1 1.2.2 1.2.3 1.3 1.3.1 1.3.2 1.3.3 1.3.4 1.4 1.4.1 1.4.2 1.4.3 1.4.4 1.4.5 1.5 1.5.1 1.5.2 1.5.3 1.5.4 1.6 1.6.1 1.6.2 1.6.3 1.6.4 1.6.5 1.7 1.7.1 1.7.1.1 1.7.1.2 1.7.1.3 1.7.2 1.7.2.1
Symbols and Abbreviations Introduction Occurrence and Processing Occurrence Mining Extraction and Refining Properties and Physical Metallurgy Chemical Properties Physical and Mechanical Properties Phase Diagrams of Tin Alloys . Strengthening Mechanisms for Tin Alloys Application of Tin and Its Alloys Pattern of Usage Pure Tin Tin Plate Tin Alloy Coatings Potential Use of Tin in Nuclear Reactors Soldering Alloys Property Requirements of Soldering Alloys Measurement of Wettability Mechanisms of Wetting Alloy Composition, Microstructure, and Properties Bearing Alloys Property Requirements of a Bearing Alloy White-Metal Bearings Aluminium-Tin Bearing Alloys Bronze Bearings Underwater Bearings Other Alloys Tin-Rich Alloys Die-Casting Alloys Pewters Fusible Alloys Tin as a Minor Additive Tin Bronzes
Materials Science and Technology Copyright © WILEY-VCH Verlag GmbH & Co KGaA. Allrightsreserved.
25 26 26 26 27 28 31 31 32 32 35 38 38 38 39 42 43 43 43 44 44 46 49 49 50 54 54 55 55 55 55 55 55 55 56
24
1.7.2.2 1.7.2.3 1.7.2.4 1.7.2.5 1.7.2.6 1.8 1.8.1 1.8.2 1.8.3 1.8.3.1 1.8.3.2 1.8.3.3 1.9
1 Tin
Tin in Dental Amalgam Superconducting Nb 3 Sn Tin in Titanium Alloys Tin in Zirconium Alloys Tin in Steels and Cast Irons Tin Compounds General Commonly Used Inorganic and Organic Compounds Applications of Compounds Inorganic Compounds Organic Compounds Organotin Compounds References
57 58 60 60 61 62 62 63 63 63 68 68 69
List of Symbols and Abbreviations
List of Symbols and Abbreviations C F G h H Hc2 Jc Tc V
atomic concentration force Gibbs free energy height enthalpy upper critical field critical current density critical temperature immersed volume
y Q 9
specific surface energy; liquid-vapor (or liquid-flux) interfacial tension specific gravity contact angle; wetting angle
b.c.c. b.c.t. BHN EPMA f.c.c. PVC ss TEM UTS VHN YS ZHS ZS
body-centred cubic body-centred tetragonal Brinell hardness number electron probe microanalysis face-centred cubic poly(vinyl chloride) solid solution transmission electron microscope ultimate tensile strength Vickers hardness number yield stress zinc hydroxystannate zinc stannate
25
26
1 Tin
1.1 Introduction Tin is one of the rarer of the common elements in the earth's crust, with a crustal abundance of about 1.5 ppm. It is as scarce as uranium and tungsten and very much scarcer than many of the common metals such as lead, zinc, and copper. Historically, it is one of the most ancient metals. When and where it was discovered is uncertain, but there is evidence to show that it has been serving mankind in different forms for well over five thousand years. Today tin is used in the form of the pure metal, alloys, and a variety of compounds. The intention is to cover all these in this short chapter. Brief coverage is also given of the occurrence, mining, and extraction practices for tin, which are very particular to this metal, and the inclusion of these presents the true story of tin.
1.2 Occurrence and Processing 1.2.1 Occurrence Tin occurs in nature in the form of oxides and also as sulphides in a variety of minerals, but it does not occur in its native state. The most common mineral amongst the oxides is tinstone or cassiterite (SnO2) and those of the sulphide-type are stannite (Cu2S • FeS • SnS2), teallite (PbSnS2), cylinderite (PbSn 4 FeSbS 14 ), and canfieldite (Al8SnS6). Of all these minerals, cassiterite accounts for most of the production of tin today. Primary lode deposits of cassiterite are believed to have been formed by hydrothermal action. From the magmatic chamber, tin was perhaps transported as gaseous tin fluoride or chloride, which on reaction with water or steam was converted into oxide, forming the mineral cassiterite. As a consequence of this, there was a release of hydrofluoric or hydrochlo-
ric gases, which reacted and altered the adjacent rocks, leading to the formation of fluorine-bearing minerals such as fluorite, topaz, muscovite, etc., thus explaining the co-occurrence of cassiterite and fluorinebearing minerals. The deposition of cassiterite occurred mostly in the form of veins in pre-existing fissures in the surrounding rocks. The mineral cassiterite is dark brown or black in colour, heavy and quite hard, it has a specific gravity of 6.8-7.1 and a Moh's hardness of 6-7. It is practically inert to most chemical solutions encountered in natural environments and therefore remains chemically unaffected during weathering. But the deposits of economic significance are mostly in the form of placers, which are formed by weathering of the primary lode deposits. Eluvial placers are formed when the weathering processes result in the removal of lighter and less inert minerals from the bedrock followed by the transport of these away from the original site, thus leaving heavier and inert mineral enriched in the form of liberated mineral particles as the residual placers at the original location. Alluvial placers, on the other hand, are formed at a location somewhat away from the primary lode deposit. In this case, the heavy and inert mineral liberated as a result of weathering is transported or carried away from the original location, for example, by a water stream. It is then deposited at another location where there is a change in the topography or the gradient of the ground. Such alluvial placers are formed at the base of the valley by the weathering of lode deposits of the hill. The sulphide minerals are also of hydrothermal origin and are usually quite complex. Sulphide-type tin deposits invariably contain a certain proportion of tin as cassiterite. The largest tin deposits occur as alluvial placers in Southeast Asia, i.e., Malaysia,
1.2 Occurrence and Processing
Thailand, Indonesia, and Mynamar. This belt is believed to extend into the Yunan province of China. Another large placer deposit is in the Amazon basin at Pitinga and Surucus in Roraina in Brazil. The recovery of tin from placer deposits is economically viable even when they contain as little as 0.025% tin oxide. Major lode-type cassiterite deposits occur in Bolivia, Australia, and Cornwall (U.K.). The deposits in Bolivia contain a significant amount of silver and are quite complex. Australian lode deposits have a substantial amount of tantalum. The deposit at Cornwall is now almost exhausted, but it has produced so far cumulatively about 5.5 million tonnes of tin. It was one of the largest suppliers of tin in the world up to the middle of the nineteenth century. At present the depth of this deposit has reached below the sea level. It may be relevant to point out here that both cassiterite and sulphide-type minerals invariably contain other valuable metals. Cassiterite is a particularly rich source of tantalum. Not only niobium and tantalum are present in the cassiterite lattice but also the minerals columbite and tantalite are associated with cassiterite. 1.2.2 Mining Mining of tin deposits depends on the type of deposits. Placer-type deposits in Southeast Asia and in Brazil are mined either by dredges or by a combination of hydraulics and gravel pumps. Mining by dredges involves the creation of an artificial pond, the deposits from the bed or from the sides are removed by buckets or suction cutters and subjected to a preliminary, rough concentration by sizing and gravity separation techniques on the dredge itself. Separated waste, i.e., sand and associated gangues is reverted back into the far side of the pond. Mining by a
27
combination of hydraulics and gravel pumps is usually adopted for mining small deposits. It involves the breakdown of tinbearing soil by powerful jets of water, the slurry is collected in a sump. The gravel pump then elevates the slurry and passes it through a sluice box or a palong (a wooden chute with a gentle slope having riffles at periodic intervals). Cassiterite, being heavier, is retained in the palong as a rough concentrate. The rough concentrate obtained by any of these two techniques is further processed in dressing plants using a variety of physical beneficiation techniques, such as conventional gravity separation by jigging and tabling, separation by modern gravity separators (such as Bartle's Mozley separator, Bartle's and Vanner's cross belt separators), separation by density using heavy media separation, and separation by using specific properties - froth flotation using the differences in the surface properties, electrostatic separators using the differences in the electrical properties, and magnetic separators based on differences in the magnetic susceptibilities. Modern gravity separators have been developed mainly to recover the finer fractions of heavy minerals from placer deposits. Many of these have been successfully utilised to improve the mineral recovery. The concentrates obtained after physical beneficiation of alluvial or eluvial deposits are relatively free of impurities and invariably contain over 70% tin. Lode-type deposits are mined by the conventional technique of open pit or underground mining. The ore is crushed and ground to facilitate the liberation of mineral particles prior to concentration. The extent of grinding required depends on the liberation size of the mineral particles. Generally, the finer the size, the higher the energy required and the lower the recovery of mineral particles in the final concen-
28
1 Tin
trate. Being fine grained, especially in veintype deposits, it is difficult to recover tin by gravity concentration. These ores are therefore subjected to flotation. Tin concentrates obtained after physical beneficiation from lode- or vein-type deposits are more complex and are usually of a lower grade (~ 20-60% tin) compared to those obtained from placer deposits. Moreover, the recovery is also lower than for placer deposits. These concentrates require further treatment prior to metallurgical smelting. The removal of sulphur and some of the impurities is accomplished by oxidising roasting and/or chloridising roasting. In recent times, the trend has been to adopt a combination of physical and chemical beneficiation techniques to obtain purer concentrates at high recoveries (Denholm et al, 1984; Halsall, 1984). The process essentially involves first the preparation of a pre-concentrate by physical beneficiation techniques with a high recovery and then treating the pre-concentrate containing up to 4-5 % tin by fuming with sulphides or chlorides to selectively volatilise tin as tin sulphide or tin chloride (Flett et al., 1984). The sulphidising process is applicable for tin values down to even 1 % or less, whereas fuming by chloridising is applicable to concentrates that contain over 5 % tin. Pilot plants and commercial plants have been set up in Australia, Bolivia, Germany, and China to treat ores/concentrates with an assay of up to about 5% tin. During the collection process itself, tin sulphide or chloride fumes are converted to oxide, which then serves as the starting material for the production of tin. 1.2.3 Extraction and Refining The industrial production of tin from cassiterite is carried out by reduction with
carbon at 1200-1300 °C. The reaction is carried out in a reverberatory furnace, or a blast furnace, or an electric furnace. Relatively pure cassiterite concentrates from placer deposits are smelted directly. Other concentrates from complex deposits usually contain some undesirable impurities and therefore require pre-treatment prior to smelting. One such treatment is roasting with or without the presence of fluxes and leaching of the roasted mass. During roasting, sulphur and arsenic are removed as oxides/sulphides. The removal of bismuth, copper, and zinc, if present, is accomplished by roasting with soda ash or sodium sulphate and water leaching of the roasted mass. The removal of antimony, bismuth, lead, and silver is generally accomplished by chloridising roasting followed by acid leaching. The concentrate with or without pretreatment is smelted in two stages to tin metal. The first stage of smelting is carried out to produce tin metal and a slag relatively rich in tin (~10% tin 1 ). Tin-rich slag is then smelted at a higher temperature in the second stage to yield impure metal and a discard slag low in tin (less than 1% tin). The impure metal is then either recirculated back to the first stage smelting or is allowed to cool down to the melting point of tin to give liquid tin and a solid iron-tin alloy (tin content ranging between 40 and 50 percent) called hardhead. Hardhead is recirculated to the first smelting stage to improve the overall tin recovery. Iron in the hardhead acts as a reducing agent during the first stage of smelting. If the starting concentrate contains too much of iron, the recirculating load for the smelting furnaces in the form of hardhead increases substantially. This has prompted 1 Weight percentages are given unless otherwise specifically indicated.
1.2 Occurrence and Processing
the development of new smelting schemes. One such scheme involves the use of single stage reduction smelting followed by slag fuming (Foo and Lightfoot, 1987). In slag fuming, the molten slag is treated with sulphur or a sulphide mineral to form volatile tin sulphide, thereby recovering tin as tin fumes which are converted to oxide and smelted. The residue is a discard slag that contains as little as 0.03-0.05% tin. Besides the conventional reverberatory, blast, and electric furnaces, a large number of new furnace options are now available for carrying out metalurgical operations. Smelting operations can be carried out in rotary furnaces, top blown rotary converters, and by the submerged combustion Sirosmelt process. Each of these newly developed processes/reactors involve the use of conditions that cause highly turbulent reactions, which result in increased specific smelting rates in contrast to the relatively slow smelting rate achievable in reverberatory furnaces (Floyd, 1980). In the Sirosmelt process, by regulating the height of the lance in the molten bath and by controlling the proportion of fuel-air mixture the reaction conditions can be controlled to produce either tin metal or tin oxide fume. Similar control of the fuel - air mixture in the top blown rotary converter can be utilised to produce either the metal or the oxide fumes. Thus these two types of reactors can be used for smelting as well as for fuming or slag cleaning operations. Besides these two, there are several other new processing techniques for slag fuming. These are cyclone fuming, slag fuming, matte fuming, and top jetting. In cyclone fuming, the finely divided sulphidic material is blown with combustion air and auxiliary fuel in a water-cooled cyclone furnace. In slag fuming, the slag is fed into a slag bed or conventional slag fuming furnace fired with oil or pulverised coal and
29
blown with additional air to burn the sulphides. In matte fuming, the charge in lump or pelletised form is fed into a matte bath maintained in a refractory-lined vessel while air to burn the sulphides is blown through the matte. In each case, the products are an iron silicate slag and a combustion gas containing sulphur dioxide, carbon monoxide, carbon dioxide, nitrogen, and steam together with sulphur and tin sulphide. Both sulphur and tin sulphide burn above the molten bath to produce more sulphur dioxide and tin oxide fume. The overall chemistry of fuming is the same in each case but the reaction paths, process dynamics, and thermal efficiencies vary. A process flow diagram summarising the various options for processing different tin sources is shown in (Fig. 1-1.) Impure tin from the primary and slag smelting stages is further refined to remove metallic impurities that have joined the metal during smelting. Refining is accomplished in two ways: by fire refining or by electrolytic refining, or by a combination of both. Fire refining consists of two steps, i.e., liquating and boiling or poling. Liquation is carried out in a sloping hearth, reverberatory-type furnace in which the metal is heated to just above the melting point of tin. Pure tin melts and trickles down the hearth, whereas the impurities that form high melting alloys and compounds remain behind. After the first run, the tin is removed and the temperature is raised further to get additional metal in the form of impure tin which can be treated further. The dross that remains is sent back to the smelting stage. Relatively pure tin obtained in the first run from liquation is further refined by poling or boiling. During this process, the liquid tin is stirred with green wood or with compressed air to cause vigorous turbulence and accomplish the oxidation of impurities to form dross.
30
1 Tin
TYPE OF DEPOSIT
SULPHIDE MINERALS
CASSITERITE
TIN SOURCE
LODE
PLACER
I
LODE
VEIN
GRAVITY SEPARATION
GRAVITY SEPARATION
VEIN
I J_
FROTH FLOTATION GANGUE DISCARD
SULPHIDE CONCENTRATE
CASSITERITE CONCENTRATE OR PRECONCENTRATE
PHYSICAL AND CHEMICAL BENEFICIATION
REDUCTION
FUME
REFINING
PRODUCT
REFINED TIN METAL
Figure 1-1. Various options for the treatment of tin sources for the extraction of tin.
The dross is then skimmed off. Fire-refined tin usually has a purity of about 99.85%. Further refining to 99.99% tin, if necessary, is accomplished by electrorefming in aqueous electrolytic cells operating with an acidic electrolyte of sulphuric acid and fire-refined tin as the anode. During the last twenty-five years or so, considerable efforts have been directed not
only at improving the existing processes but also at developing alternative processing schemes involving the chlorination and partial reduction of cassiterite. Tin from chlorides and partially reduced cassiterite (Pearson et a l , 1977 a, b) has been recovered in a variety of ways. Most of the alternative processes have been developed to laboratory scale or they have reached the
1.3 Properties and Physical Metallurgy
stage of bench level or pilot plant-scale tests. In addition to the primary sources, some quantity of tin is also recovered from its secondary sources such as industrial wastes and scraps. Lightly contaminated scrap of high tin alloys is usually recycled by melting, refining with respect to harmful association, and finally adjusting the composition for reuse in the same or other applications. Heavily contaminated scrap of high tin alloys is usually treated by a sequence of chemical and metallurgical operations to purify the metal with respect to all the impurities, including all the alloying constituents and contaminating impurities. Thus the demand on the purification steps in recycling is much more than is imposed during the treatment of primary sources. Another source of tin is tin-plate scrap and used tin cans. Tin from these sources can be selectively dissolved or stripped in many ways, such as chemical or electrolytic dissolution in caustic soda solution in the presence of oxidising agents or selective chlorination of tin with dry chlorine at about 50-80 °C (Germain, 1983). In the first two cases, tin is obtained as sodium stannate in aqueous solution and in the last one as tin chloride. Tin from all these is then recovered as an oxide, any other compound, or as tin metal.
1.3 Properties and Physical Metallurgy 1.3.1 Chemical Properties
In the periodic table tin occupies the same group as carbon, i.e., the fourth group and therefore it has some of the attributes associated with carbon. One of the allotropes of tin, i.e., grey tin, has the diamond cubic structure, high resistivity, and
31
semiconducting characteristics. Tin also has a tendency to form tin-carbon organotin chain compounds. In tin, the valence shell or the outermost electronic shell has four electrons, two each in the s and the p orbitials of the fifth orbit, thus exhibiting two stable valency states of 2 and 4. Chemically, tin is amphoteric in nature, i.e., it reacts with strong acids and alkalies but is resistant to nearly neutral solutions, and it therefore has good corrosion resistance in many weak acids, alkalies, and neutral solutions. In aqueous solution the presence of oxygen and/or oxidising salts may greatly increase the corrosion rate. Tin does not react with many of the gases such as nitrogen, hydrogen, carbon dioxide, or gaseous ammonia but it reacts readily with the halogens and moist sulphur dioxide, for example. Tin exhibits a very interesting electrochemical behaviour with respect to iron. In normal environmental conditions, tin is cathodic with respect to iron. A coating of tin on iron therefore protects the iron by anodic protection. If, however, there are pin holes in the coating enhanced corrosion and pitting of the iron base will occur. In acidic media in the absence of air, there is a reversal of the potential in the iron-tin electrochemical couple, due to which tin becomes anodic with respect to iron. It is this property of reversal in the electrochemical behaviour that is utilised for making tin cans for food and beverages. As the tin is anodic in tin cans, there is preferential corrosion of the tin. Tin salts taken up by the solution inhibit the corrosion of iron, thus protecting the food from the adverse effect of iron contamination. Tin does not affect the flavour and the taste of the food, and it does not discolour it. Most importantly, a certain amount of tin is acceptable in food items because of its nontoxic nature.
32
1 Tin
1.3.2 Physical and Mechanical Properties Tin has ten naturally occurring stable isotopes with an average atomic weight of 118.69. It has two allotropic forms, white and grey tin with an allotropic transformation temperature of 13 °C. The low temperature form, a, is termed grey tin, it is semiconducting in nature, has high resistivity, and a diamond cubic structure. The structure of white tin (P) is body-centred tetragonal. It is soft, ductile, and metalic in character. The volume difference between the two allotropes is about 27 %. Therefore on transformation from white to grey tin, considerable swelling occurs, which may appear as hemispherical protuberances or warts on the surface. This is referred to as tin pest. The transformation from P to a is very sluggish at 13 °C, but is quite rapid at -40 °C. The allotropic transformation and its rate depend on the temperature, purity level, and previous mechanical history. Impurities such as lead, antimony, and bismuth are invariably present in the commercial grade metal. All of these impurities delay the onset of transformation from a few hours up to several hundred days depending on the impurity content. It has also been observed that even for pure tin the transformation from P to a does not occur after three years of exposure to -40 °C if the metal does not have any preexisting nuclei. Nucleation can be induced by the presence of heterogeneities such as second phase particles having a diamond cubic structure or strain centres produced by repeated transformation cycles. Another peculiar characteristic associated with pure tin is 'tin cry' or the cracking noise emitted during plastic deformation. This is particularly noticeable during the bending of cast, coarse-grained sticks and is due to the twinning of the crystals on {301} planes. Some of the selected physical and
mechanical properties of tin are presented in Table 1-1. 1.3.3 Phase Diagrams of Tin Alloys Tin has a strong tendency to form ordered intermetallic compounds with a number of elements. This is reflected by the occurrence of several intermetallic compounds with narrow composition ranges in a number of equilibrium phase diagrams of tin. Some of the important intermetallic compounds of tin are listed below: a) Cubic LI 2 structure (Cu3Au type): CaSn 3 , LaSn 3 , CeSn 3 , PrSn 3 , USn 3 , Pd 3 Sn, and In3Sn. b) Hexagonal D0 1 9 structure (Mg3Cd type): Ti3Sn, Mn 3 Sn, Fe3Sn, Ni3Sn, and Pt3Sn. c) Hexagonal BSX structure (NiAs type): Mn 3 Sn 2 , Fe 3 Sn 2 , Co 3 Sn 2 , Ni 3 Sn 3 , Rh 3 Sn 2 , and Pd 3 Sn 2 . d) Tetragonal C l 6 structure (Cu2Al type): MnSn 2 . FeSn 2 , CoSn 2 , and RhSn 2 (>500°C). Depending on the nature of the phase reactions present, the phase diagrams of tin can be broadly classified into the following four classes: eutectic phase diagrams, peritectic phase diagrams, phase diagrams showing a number of intermetallic compounds, and phase diagrams showing liquid phase immiscibility. Binary phase diagrams of tin with bismuth, gallium, lead, thallium, or zinc are of the simple eutectic type. The presence of low melting eutectics in these alloys makes them suitable for applications as fusible alloys. A simple eutectic diagram of the tin-lead system is shown in Fig. 1-2 alongside a tin-indium diagram in which the eutectic forms between two intermediate phases P and y which have pseudobody-centred tetragonal and hexagonal
1.3 Properties and Physical Metallurgy
33
Table 1-1. Physical and mechanical properties of tin. Property Atomic no. Atomic weight Crystal structure Density
Conditions
p tin (white tin) a tin (grey tin below 13 °C) P tin at 15 °C a tin at 15 °C liquid tin
Melting temperature Boiling point Thermal conductivity Coefficient of linear thermal expansion Specific heat Electrical resistance
Superconductivity transition temperature Electron work function Young's modulus Tensile strength
Brinell hardness
structures, respectively. A plot (Fig. 1-3) of the liquidus temperature as a function of the atomic fraction of the alloying elements (zinc, lead, thallium, cadmium, bismuth, indium, and gallium) shows that the liquidi of the tin-rich side of these eutectic diagrams nearly coincide. Though the liquidi corresponding to the alloy-rich side are not linear, the eutectic temperatures in these systems follow the same order as the melting points of the respective alloying elements. The phase diagram of tin with antimony, silver, and niobium are characterised by two or more peritectic reactions. As the liquidus drops from the high melting metal to the tin side, these peritectic reactions are encountered. The tin-antimony phase diagram (Fig. 1-4) shows the presence of the (3-SbSn phase, which is an important mi-
ptin at 100 °C 50 °C a phase at 10 °C P phase at 25 °C ptin at 20 °C P tin at 200 °C a tin at0°C
at 20 °C at20°C atl00°C at200°C atl7°C
Value 50 118.69 b.c.t. f.c.c. 7.29 g/cm3 5.77 g/cm3 6.976 g/cm3 231.88 °C 2625.0 °C 60.7 W/(m K) 23.1 um/(m K) 205.0 J/(kg K) 222.0 J/kg K) 12.6 \iQ cm 23.0 uD cm 300.0 jiiQ cm
3.73 K 4.64 eV 49.9xlO 9 N/m 2 14.5xlO 6 N/m 2 11.0xl0 6 N/m 2 4.5 x 106 N/m2 4.02 x 106 kg/m2
croconstituent of a large number of bearing alloys. In both eutectic and peritectic systems the terminal solid solutions generally exhibit solubility of the alloying element to a limited extent. The phase diagrams of tin with low melting alkali metals like potassium and lithium show a number of line compounds either congruently melting or forming through a peritectic reaction at temperatures much higher than the melting temperatures of the constituent elements. The liquidi in these cases show a maximum at an intermediate composition level (Fig. 1-5). The phase diagrams of tin-selenium and tin-sulphur (Fig. 1-6) are similar, showing a steep rise in the liquidus temperature from the melting temperatures of the two pure components. The solubility of one component in the terminal solid solu-
34
1 Tin Weight Percent Lead 50 60 70 80
(a) 0 10 20 30 400 r-T 1 1
2.0
40 r
30
40
50
60
70
80
90
80
90
Atomic Percent Lead
(b)
10
20
30
Weight Percent Indium 40 50 60 70
Pb
23I.9 224°C 200 I56.6 150
100 (In)
30
40
50
I
l
I
i
60
70
80
90
Atomic Percent Indium
tion of the other is extremely small. Even the liquid phases in these systems do not mix, as represented in phase diagrams by phase fields containing two liquid phases. Tin goes into solid solution to a significant extent in alloys based on zirconium, titanium, hafnium, nickel, palladium, or magnesium, and is therefore added as an alloying element in many of the commercial alloys based on these metals. The solubility of these elements in tin is, however,
Figure 1-2. (a) The tin-lead and (b) the tin-indium phase diagrams (Massalski, 1986). I00 n
extremely small. The phase diagram of all these systems can be grouped into one class, which is characterised by a steeply increasing liquidus from the tin side to the congruent melting point of an intermetallic which in turn forms a eutectic with the terminal solid solution. The presence of several other intermetallics involving peritectic and peritectoid reactions makes these phase diagrams more complex. The phase diagram of tin-zirconium is shown in Fig. 1-7 as a protoype.
1.3 Properties and Physical Metallurgy
400 -
"I00 90
80
70
60
50
40
30
20
10
0
ATOMIC %TIN
Figure 1-3. A plot showing liquidus lines for various binary eutectic phase diagrams of tin.
600 -
Sn
10 20
30 40 50 60 70 80 WEIGHT PERCENT ANTIMONY
90
Sb
Figure 1-4. The tin-antimony phase diagram (Massalski, 1986).
1.3.4 Strengthening Mechanisms for Tin Alloys
Pure tin is a very soft metal, its ultimate tensile strength at room temperature being slightly over 106 Pa. For a variety of applications, tin has to be strengthened using the available techniques of alloy hardening. Solid solution strengthening is often employed but the scope of this is rather restricted due to the limited solubility of only a few elements in tin. The solubility limits
35
of zinc, cadmium, and bismuth are below 2 % while those of antimony and indium are in the range of 6-7%. These elements are primarily used as alloying additions for strengthening tin alloys. The effectiveness of this solid solution strengthening, as shown in Fig. 1-8 (Hedges et al., 1960), reaches a maximum for zinc, which has the lowest solubility of the elements mentioned. This is not unexpected as a larger* difference in the size of the solute and the solvent atoms tends to restrict the extent of solubility while increasing the efficacy of strengthening. A combination of two or more solutes is known to be more effective, as seen in alloys containing about 1 % cadmium and 3-9% antimony. Strengthening by the distribution of a second phase is commonly employed in a number of tin alloys. Ordered intermetallic compounds such as Ni 3 Sn 4 , FeSn 2 , Cu 6 Sn 5 , and SbSn dispersed in the matrix of a tin-rich solid solution can cause significant hardening of inherently soft tin. The extent of coherence between the precipitate and the matrix for different combinations has not been studied in detail. The fact that dispersed particles of alumina and silicon carbide are not capable of hardening tin alloys suggests that coherence of the precipitate phase, at least to a limited extent, is a necessary factor for strenghtening. Strenghtening through the precipitation of intermetallic compounds can be achieved by the conventional age hardening treatment involving solutionising followed by ageing. The alloys that are most amenable for hardening by ageing treatments contain cadmium (1 -2%) and antimony (9-14%). Chill casting of these alloys results in the formation of supersaturated solid solutions which on ageing in the temperature range of 100-140°C produces hardening due to the distribution of SbSn or CdSb precipitates.
36
1 Tin Weight Percent Lithium 2
100 0 Sn
5
10
20 30 50 I00
Figure 1-5. The tin-lithium phase diagram (Massalski, 1986). 10
20
30 40 50 60 70 Atomic Percent Lithium
Weight Percent Sulphur 10 20 30 40
80
60
90
I00 Li
80 100
Figure 1-6. The tin-sulphur phase diagram (Massalski, 1986). Atomic Percent Sulphur
Some of the intermediate phases on the tin-rich side of binary or multicomponent phase diagrams have high hardness. For example, the (3-alloys (A3 structure) of the cadmium-tin system have a hardness higher than 30 VHN in the quenched condition. However, the p-phase decomposes into a mixture of tin-rich and cadmiun-rich phases due to room temperature ageing, and this is accompanied by a gradual softening. The y-phase (hexagonal structure) of the tin-indium system is stable at room
temperature but does not possess very high hardness ( - 1 0 - 1 2 VHN). Since the (3and the y-phases are isomorphous, it is possible to make ternary alloys based on the tin-cadmium-indium intermediate phase with high hardness and stability at room temperature. The commonly used strengthening method of work hardening is not suitable for tin alloys. This is due to the fact that the hardness of pure tin and some of its alloys decreases with increasing cold work
37
1.3 Properties and Physical Metallurgy
IO
Weight Percent Zirconium 30 40 50 60 70
20
80
90
I00
400 23I.9°C X
Figure 1-7. The tin-zirconium phase diagram (Massalski, 1986).
200
0 Sn
20
30
40
50
60
70
80
Atomic Percent Zirconium
(Fig. 1 -9) when the latter exceeds a limit of about 20% (Hedges et al., 1960). This anomalous behaviour can be explained in terms of the spontaneous recrystallisation of tin-based alloys at room temperature. When the cold work exceeds the critical
90
I00 Zr
value for nucleation of recrystallised grains, spontaneous recrystallisation results. This explains the phenomenon of work softening in these alloys. It may be noted that the hardness drops to a value lower than the original with increasing
6%Sb,2%Cu,2%Bi
3
4
5
WT.% SOLUTE
Figure 1-8. The effect of solute content on the hardness of tin-based solid solutions (Hedges et al. 1960).
10 20
30
40
50
60
70
80 90
100
% REDUCTION BY ROLLING
Figure 1-9. The effect of cold rolling on the hardness of two tin alloys (alloy compositions are indicated against the respective plots) (Hedges et al., 1960).
38
1 Tin
cold work. This may be due to the removal of residual stresses possibly present in the starting material. In addition, the development of a favourable crystallographic texture can contribute towards the softening process.
1.4 Application of Tin and Its Alloys 1.4.1 Pattern of Usage The first recorded use of tin dates back to about 3000 B.C. for making bronze weapons and implements and also for making ornaments. Since then many new and diverse applications have been discovered. Tin as a metal is almost always used in association with other metals, either as a coating or as an alloying constituent. This is because its low mechanical strength prevents the use of pure tin as a structural component. In present day industry, tin however owes its place of pride to certain other basic properties of the metal, i.e. its low melting point, nontoxic nature, chemically less reactive nature, and tendency to wet most of the metals. The low melting point and good wetting properties of this metal and its alloys account for their use as solders. Non-toxicity and resistance to corrosion are responsible for its used in tin plate for the packaging of food and beverages. Its ability to form alloys with many common metals is utilised in making various commercial alloys. The metal has an ability to retain oil film and this makes it ideally suited for use in bearing alloys. In addition to these general applications, tin also finds specific applications in many other areas such as nuclear energy, aerospace, superconducting alloys, in the manufacture of glass, and in the preparation of organic and inorganic compounds of tin.
The total annual production of tin in the world has been in the range of 160000 200000 tonnes for the last two decades. The major use of tin until about 1984 was in the tin plate industry (Carlin, 1985; Pitman, 1988). Advancements in electronic and automotive industries have led to a substantial increase in the application of tin in solder alloys and at present this accounts for the largest proportion of tin consumption. Some data compiled on the, production and the consumption pattern in selected countries are presented in Table 1-2. Some of the important features of various applications of tin, tin alloys and compounds are described in the following sections. 1.4.2 Pure Tin Pure metallic tin is used in a few very specific applications. As tin powder, it is used in making bronze parts and as a sintering aid in making sintered iron parts, it is also used as a gun lubricant and as an additive in time delay devices in gun charges. The powder is industrially produced by the atomisation of liquid metal with air. The atomised tin powder particles, as illustrated in Fig. 1-10, have a
Table 1-2. Pattern of tin usage in selected countries (in percent). Application Tin plate Solder White metal Bronze and brass Chemical Others Total tin consumption (tonnes)
1984
1986
31.8 29.1 6.3 6.4 10.0 16.4 153 500
29.4 29.7 — — 13.9 26.9a 148 700
Including white metal and bronze.
1.4 Application of Tin and Its Alloys
Figure 1-10. The morphology of atomised tin powder particles.
spherical morphology and a fairly uniform size. As tin foil, it is used for making electrical capacitors and laminated tin foils for lining bottle caps. As wire, tin is used in making fuses and safety plugs. Tin tubes, pipes, and tin-lined containers are used for conveying and/or storage of distilled water, carbonated beverages, beer, and wine. Tin in the form of a molten bath is used for making glass sheets by Pilkington's float glass process (see Sec. 1.4.4.3 in Volume 9 of this Series). In the float glass process, molten tin provides an optically flat surface on which molten glass is cast, thereby eliminating the cumbersome polishing step. The glass sheet thus produced are suitable for use as window panes, mirrors, automobile windshields, etc. Pure tin is also used for making anodes for the electrolytic deposition of tin for manufacturing tin plate and tin alloy coatings. 1.4.3 Tin Plate Nearly one third of the total tin produced in the world today goes into making tin plate, which is in fact a thin sheet or strip of low carbon steel (containing about 0.1 %C) coated on both surfaces with a thin layer of tin. Tin plate has the desirable features of strength, durability, fabricabili-
39
ty of steel, good corrosion resistance, compatibility for storage of many of the chemicals and food articles, and the aesthetic appearance of tin. Most of the pre-coated sheets are then utilised for making tin cans for food and beverages, and containers for paints, aerosols, and many other articles. About ten percent of the total tin plate produced is used in applications other than packaging, which include the automotive, electrical, and electronic industries and for a variety of decorative articles. In such applications corrosion resistance, solderability, and aesthetic appearance are important. Since most tin-plate applications are dependent on the surface area, i.e., how many pieces can be made from a given area, tin plate is traded on the basis of area and not weight. The normal unit is a SITA (System International Tinplate Area) of 100 m 2 . The thickness of the tin coating on tin plate may vary from less than a micrometre to a few micrometres and may either be the same on both the sides or different. The thickness of the tin coating on tin plate is usually expressed in terms of g/m2, for example a thickness of 11.2 g/m2 corresponds to 1.54 jam. The thickness of steel sheets or strips normally used for making tin plate is between 0.15 and 0.5 mm. The mechanical properties of tin plate essentially depend on the properties of the steel sheet, which in turn can be controlled by controlling the rolling and annealing procedure adopted during the fabrication schedule. Since the fabrication and coating processes are independent, it is possible to produce tin plate with a desired combination of mechanical properties and coating thickness. Tin plate in commercial practice is designated by the coating thickness, the chemical composition of the steel, and the temper. The product having an equal thickness of tin on both the sides is desig-
40
1 Tin
nated by the letter E and that having different thickness by D. For example, E 5.6/5.6 denotes tin plate which has a thickness of 5.6 g/m2 of tin on each side and D 11.2/2.8 denotes tin plate which has a thickness of 11.2 g/m2 of tin on one side and 2.8 g/m2 on the other side. Most of the steel used for manufacturing tin plate is currently produced by the oxygen steel making process. The chemical composition is controlled within a specific range by the appropriate selection of raw materials and process parameters. Steels used for tin plate generally contain carbon in the range of 0.04-1.12%, sulphur 0.015-0.05%, phosphorus 0.015-0.02%, silicon 0.01-0.02%, copper 0.02-0.2%, and manganese 0.2-0.6 %. Generally nickel, chromium, and molybdenum are not deliberately added, but they may be present in the range of 0.04-0.06 %. To achieve improved corrosion resistance for packaging certain types of food, it is necessary to limit the content of residual elements in the steel each to a value of less than 0.02%. The steel ingots of desired chemistry are rolled by a combination of hot and cold rolling to form strip or sheets of the desired thickness. The rolling is generally carried out continuously in a tandem mill. After hot rolling the strip is descaled, cleaned, and cold rolled imparting a cold reduction of 80-90%. The sheets are then cleaned once again and annealed in a nonoxidising atmosphere. Appropriate control of the extent of cold reduction and the annealing cycle provides the desired combination of mechanical properties with a UTS of approximately 330-670 MPa and a varying degree of ductility, stiffness, and directional properties. Annealing is carried out mostly below the ferrite to austenite phase transformation temperature to achieve recrystallisation and grain growth. The steel
sheet is then either subjected to temper rolling with a cold reduction of 1 - 4 % to achieve the final properties and surface finish, or it is subjected to another coldrolling cycle to produce sheets for manufacturing double-reduced tin plate, which is stiff and strong. The double-reduced sheets, however, have very highly directional properties. The steel sheets are inspected for thickneses, any surface defects, pinholes, etc. and the defective material is removed; the edges are trimmed from the sheets, which are cleaned by degreasing and pickling. These are then ready for the application of the surface coating of tin. A tin coating on steel sheets is achieved either by dipping the steel sheet in a molten bath of tin, termed "hot tinning" or by electrolytic deposition of tin from aqueous solutions. In the world, most of the production (nearly 98 % or more) of tin plate is by electrolytic deposition. The electrolytic process is suitable for applying a very thin and uniform coating of tin, it also offers precise control of the thickness of the coating, moreover it is highly suited for applying a different coating thickness on either side of the sheet. Electrolytic tin plate is produced by the continuous electrodeposition of tin onto steel strip moving at a high speed (about 10 m/s) through a processing line with uncoilers, shear, and a welding and looping arrangement on one end and coilers and shear on the other. The strip initially passes through a cleaning line comprising alkaline solutions to remove oil and grease, followed by water rinsing. Cleaned strip then moves through an electrodeposition line containing aqueous tin solution as an electrolyte and tin anodes; the strip acting as the cathode. Spent tin anodes are removed and replaced periodically. The electrolytes for manufacturing tin plate are proprietary but they can be
1.4 Application of Tin and Its Alloys
broadly classified into four different categories, i.e., ferrostan (based on a combination of stannous sulphate and phenolsulphonic acid), halogen (based on stannous chloride/fluoride), acidic stannous fluoborate, and alkaline stannate. Of these, only the alkaline stannate electrolyte produces a smooth tin deposit without any additives, but the deposition rate is half that of the other electrolytes. All the other electrolyte require some organic additives to obtain a smooth and coherent deposit. The deposition line may either be laid out horizontally or vertically. A horizontal line is used for the halogen process. The anode in this process is placed horizontally and the strip moves over it, thereby resulting in a coating on one side, i.e., the facing side of the strip in one series of tanks. The other side is coated in another series of tanks. Vertical electrolytic lines are used for all the other electrolytes, in these lines the anode is hung vertically between two strips, thus tin plating is simultaneously accomplished on both sides. As plated tin has a dull, matt appearance, the physical appearances and the functional performance of tin plate can both be improved by a flow brightening treatment. In this treatment, the strip is electrically heated for a very short time to slightly above the melting point of tin by resistive or induction heating. The strip is then plunged in water. During this treatment the surface is brightened due to surface melting, also there is an interaction between the coating and the steel sheet which results in the formation of a thin layer of the FeSn2 intermetallic compound at the interface. The formation of this compound not only improves adhesion but also the corrosion resistance of the tin coating. Flow-brightened sheet is further protected against excessive oxidation and rough handling by first a passivation treat-
41
ment in a chromate bath and then coating with an oil film on the surface. Tin anodes used for electrolytic deposition generally have a purity of 99.80 %, but there is an increasing trend to use a higher purity metal. The coating thickness of tin in electrolytic tin plate is usually in the range of 0.3-0.8 ^im. Now with a better understanding of the effect of different impurities and the availabily of higher purity tin (with good control of the harmful impurities), it is possible to reduce the coating thickness to much below 0.1 |im without adversely affecting the quality of the tin plate. Another process for manufacturing tin plate is by hot tinning, this involves immersing the steel sheet or strip in a molten bath of tin. It can be accomplished either in a batch process or in a continuous process. The continuous process involves passing cleaned and pickled sheet first through a molten flux (which may be zinc chloride or a mixture of zinc and ammonium chloride), then through a molten tin bath maintained at a temperature between 280 and 325 °C, and finally through a grease pot or a palm oil bath. All these are generally integrated into a compact, single assembly. The thickness of the tin coating is controlled by appropriate selection of the temperature of the bath, the retention time of the sheet, and the pressure on the spring-loaded rolls placed in the grease pot. In fact, the rolls in the grease pot scrape off excess tin on the tin plate and facilitate better control of the thickness. The tin plate produced by hot tinning is generally thicker and has a thicker interlayer of FeSn2 compared to that obtained by the electrolytic process. The nature of the coating depends on the amount and type of impurities present in the tin. Since the as-produced coating is bright, it does not require any additional treatment. The
42
1 Tin
post-coating passivation and lubrication treatment are, however, similar to those adopted in manufacturing electrolytic tin plate. Hot tinning has also been adopted for applying a tin coating on a variety of complicated shapes. This can be accomplished by wiping or by applying tin in the form of a stick or powder onto a cleaned and preheated surface. Although a tin coating is useful for making tin plates, coatings of certain tin alloys are also useful for a variety of applications.
1.4.4 Tin Alloy Coatings Tin alloy coatings in general have a high hardness, and may be brighter and have better corrosion resistance than pure tin. Some of the alloy coatings are those of tin-lead, tin-nickel, tin-zinc, tin-cadmium, and tin-copper alloys. Of these, only tinlead can be applied either by hot-dip coating or by the electrolytic co-deposition process, whereas all the other alloy coatings are produced only by co-deposition of the alloying constituents from aqueous electrolytes. Some of the important features of various tin alloy coatings are shown in Table 1-3.
Table 1-3. Important features and properties of tin alloy coatings. Alloy
Composition range
Properties
Application areas
Coating process
SnPb
Solder composition - 60% Sn
Suitable for soldering
Fluoroborate and nonfluoroborate proprietary electrolytes
Terneplate 8-25% Sn
Corrosion resistance, solderability, improved drawability
SnNi
65% Sn
Hard, tarnish and corrosion resistant, oil retention
SnZn
70-85% Sn
Readily solderable
Electrical connectors, electronics industry for printed circuits, motor car radiators, in printing Automotive industry for fuel tanks, fuel air filters, radiators, heat exchangers, electrical conduit bases, connectors, radiators, outdoor fixtures Watch parts, piston and brake mechanisms, drawing and scientific instruments, electrical equipment, decorative light fittings, jewellery Electrical and electronic equipment, tools, brake mechanisms, hydraulic machine components, automotive applications
SnCd
20-50% Cd
SnCu
7-20% Sn
Good protection for steel in wet and marine environments High wear and Hydraulic equipment corrosion resistance Tarnish resisDomestic articles, tableware tant
Sodium stannate and cadmium cyanide, fluoride-fluorosilicate solutions Sodium or potassium stannate, copper cyanide
40% Sn
Hot-dip coating in continuous mills
SnCl2, 2H 2 O, NiCl 2 , NH 4 HF, 70°C, p H 2 - 4 , cathode current density 260 A/m2 Sodium or potassium stannate, zinc cyanide
Sodium or potassium stannate, copper cyanide
1.5 Soldering Alloys
43
1.4.5 Potential Use of Tin in Nuclear Reactors
1.5.1 Property Requirements of Soldering Alloys
A new type of nuclear reactor concept was evolved during the late 1960s which was based on the temperature dependence of the solubility of nitrogen in tin and the reactivity of uranium with nitrogen. Tin is used because of its ability to dissolve nitrogen at temperatures in the range of 14751800°C, thus making it available as a chemical reactant. In this reactor, uranium is dissolved in tin which is held under a slight nitrogen pressure in a graphite-lined reactor vessel. The uranium reacts with the dissolved nitrogen and forms a uranium nitride precipitate which sinks to the bottom of the vessel. When the mass of the uranium nitride becomes critical, nuclear fission is initiated. If, however, the temperature rises excessively, the uranium nitride dissociates and uranium goes into solution, thus reducing the amount of uranium nitride and thereby slowing down the fission chain reaction. Conversely, when the temperature drops, the rate of nitriding increases, which in turn increases the mass of the uranium nitride so increasing the rate of the fission chain reaction.This type of reactor is still far from a commercial reality.
The major requirements of a soldered connection can be listed as follows although all of these attributes are often not required simultaneously in a specific service condition: i) To provide an electrical and thermal conducting path across the joint, ii) To connect the components mechanically and to retain adequate strength at temperatures from cryogenic levels to about 220 °C. iii) To form liquid- or gas-tight seals. The selection of the composition of a solder alloy and the flux, and the choice of the soldering technique are dictated by the end use and the service condition of the soldered joints. In the field of electronics, the electrical connectivity is obviously the most important requirement. The electrical conductivities of solder alloys are relatively low, only 8-15 % of that of copper. This, however, does not introduce significant electrical resistance in the circuit because of the short conducting path and the large surface area of contact. The strength aspect of the soldered joint assumes importance where the joint is expected to withstand stresses due to gas or liquid pressure or thermal fluctuations. The quality of a soldered joint is influenced by the following three factors: the nature of the surfaces to be joined, the solder alloy to be used, and the choice of the flux which assists in obtaining intimate contact between the two workpieces. All these factors control the process of wetting of the solid base metal surfaces by molten solders. It is therefore necessary to understand the mechanism of wetting and its measurements. These aspects are discussed in the following sections.
1.5 Soldering Alloys Soldered joints are often intended to serve as a means of providing continuity for the conduction of heat or electricity, or as seals for liquids and gases. In these applications the soldered joints are not subjected to the high mechanical stresses encountered in certain other applications such as aerosol containers, motor car radiator header tanks, and refrigeration appliances.
44
1 Tin
1.5.2 Measurement of Wettability
illary forces. F can be written as
The wetting and spreading of solder on a surface may be explained in terms of the balance of interfacial tension, as illustrated in the well-known triangle of force diagram (Fig. 1-11) showing the surface tensions exerted by the interfaces between the substrate metal (M) and the flux (F), the molten solder (S) and the substrate metal (M), and that between the molten solder (S) and the flux (F). The specific surface energies, y, of the respective interfaces being denoted by yMF, ySM, and ySF respectively, the balance of force at the tri-junction can be expressed as
F=ylcos9-gv
(l-l)
SM
where 9 is the wetting angle. As the contact angle 9 decreases, better wetting is achieved and, as a result, the total surface energy decreases through the spread of low y material (solder) over the substrate material with a high y. A quantitative evaluation of wettability involves the measurement of the force to which a wire sample is subjected when partially immersed in a molten solder bath (Manko, 1964). A sensitive wetting balance, known as a meniscograph, plots the force versus time graph which gives a "history of the wetting". The force, F, experienced by a sample is the algebric sum of the buoyancy and the resultant of the cap-
r
Morten solder /flux Flux Base metal / flux
y
Molten solder / base metal
Base metal
Figure 1-11. Diagramatic representation of the interfacial tensions between the base metal, solder pool, and flux cover.
(1-2)
where F= force measured, y = liquid-vapour (or liquid-flux) interfacial tension, /=perimeter length, 9 = contact angle, Q = specific gravity of the solder at the measuring temperature, and v = immersed volume of the test sample. Since the value of y is usually unknown, a meniscometer is used for measuring the height to which the meniscus rises in a given time interval. The height h is related to y by the following relation
= {2y(l-sm9)/Q}x/2
(1-3)
In conventional practice, measurements of both force and height are taken after a time interval of 3 s in two independent experiments and the contact angle is computed using Eqs. (1-2) and (1-3). Since two separate experiments are needed for the measurement of 9, surface conditions of the samples used must be kept identical for the two cases. 1.5.3 Mechanisms of Wetting
The process of wetting is motivated by the reduction of the total surface energy of a system. Generally high melting, hard solids such as metals, metal oxides, nitrides, and glasses are associated with high surface energies in the range of 0.5-5 J/m2, while low melting soft solids usually have low surface energies below 0.5 J/m2. It is generally believed that the soldering process only involves wetting in the physical sense by virtue of van der Waals forces. However, recent work has established beyond any doubt that the formation of solid solutions or intermetallic compounds invariably occurs within the soldered layer, and that metallurgical bonding is established. It is also seen that liquid metals
45
1.5 Soldering Alloys
with the best wetting properties readily alloy with the substrate metal. Tin, the active constituent in most soft solders, both wets and alloys with many base metals, such as copper, iron, and nickel. The solid solubilities of these metals in tin are very small, and therefore very thin layers of intermetallic compounds form in the soldered joints. Electron probe microanalysis (EPMA) has shown that the reaction between solid copper and liquid tin at 923 K results in the formation of the intermetallic compounds fi and 7 which contain 23-27 at.% Sn (Hasouna et al., 1988). It has also been demonstrated that the contact angle of liquid tin on solid copper changes with time, as shown in Fig. 1-12. The progressive interaction between copper and tin with the formation of intermetallic compounds is responsible for this increasing tendency for wetting with time. The variation in the contact angle with the temperature of interaction is shown in Fig. 1-13, where a minimum is seen at a temperature of about 690 K, which is where the difference between the liquidus composition and the solidus composition is the smallest. EPMA has also established the presence of different intermetallic compounds depending on the temperature of interaction. In the temperature range of 505-703 K, the r\ phase (Cu6Sn5) has been found to form either as a layer or as finely distributed particles at the interface. In contrast, at about 790 K the formation of the 5-phase (y-brass, D8 2 structure) occurs at the interface, and at still higher temperatures ( - 9 2 3 K) a thin layer forms presumably due to the surface diffusion of tin. The atmosphere plays an important role in the wetting process. This is demonstrated by the fact that the contact angle can be significantly reduced by lowering the oxygen partial pressure below the equilibrium
25
L K
o D
CD
O
863 898 913 923
K K K K
15°
2 5
0.5
I TIME/ks
Figure 1-12. The variation of the contact angle between molten tin and solid copper with time at different temperatures in a hydrogen atmosphere (Hasouna et al., 1988).
• Hasouna et al
_o>50
l\ Yokota et al
"5 3OC §20°
° o505
VI.,
600
700
800
900
„
I.
IOOO
Temp. K Figure 1-13. The temperature dependence of the contact angle between molten tin and solid copper (reaction times are indicated on the plots) (Hasouna et al., 1988; Yokota et al., 1980).
oxygen partial pressure for Cu 2 O and SnO2 at a given temperature. The oxide barrier to the tinning of metals in many cases can be removed by reducing the oxide layer using hydrogen at high temperatures, but "fluxing", being a much more convenient method, is widely used for this purpose. The major constituent of the flux is zinc chloride which in the presence of water produces some free hydrochloric acid, as shown in the reaction ZnCl 2 + H 2 O -* Zn(OH)Cl + HC1
(1-4)
46
1 Tin
OXIDISEDSURFACE
TIN BASE METAL COMPOUND
In the case of the soldering of copper, the oxide layer is dissolved as chloride, leaving bare copper on which the molten tin gradually spreads. This process is schematically illustrated in Fig. 1-14. During fluxing the liberated hydrochloric acid reacts with the tin and the stannous chloride produced is carried in the flux. Just ahead of the advancing front of molten tin, the stannous chloride reacts with the freshly exposed copper surface according to the reaction SnCl2 + Cu = Sn + CuCL
(1-5)
Tin resulting from this reaction gets deposited on the copper surface, alloys with the base metal, and merges with the advancing, molten tin front. The dual actions of the flux in the oxide removal and in the deposition of the tin greatly facilitate the wetting of the metal surface by the molten tin. 1.5.4 Alloy Composition, Microstructure, and Properties
For the majority of soldering operations in the electronics industry, binary lead-tin alloys are chosen. The eutectic temperature being 183 °C, these alloys offer a low soldering temperature to avoid the risk of damage to temperature-sensitive components. The temperature difference between the liquidus and the eutectic temperature,
Figure 1-14. Schematic diagram showing the displacement of flux by molten solder.
often known as the "pasty gap" can be controlled by selecting alloy compositions deviating from the exact eutectic (61.9% Sn) composition. For mass soldering of printed circuit boards the 60 % Sn-40 % Pb alloy is often used, while a 50 % Sn alloy finds extensive applications in less exacting, hand soldering operations. The general microstructure of the as-cast solders of these compositions, which are on the lead-rich side of the Pb-Sn eutectic, comprise primary lead dendrites in a eutectic matrix (Fig. 1-15 a). As the solid solubility of tin in lead is sharply reduced on going from 183°C to ambient, precipitation of the tin-rich phase occurs within the primary lead dendrites. Such a precipitation results in softening of the binary Pb-Sn solders and the effect is pronounced in Pb-Sn-Sb solders where precipitation of the SbSn intermetallic compound occurs. The microstructures of solder alloys, tin and lead being their major constituents, are primarily governed by the binary leadtin equilibrium diagram (Fig. 1-2 a). Nonequilibrium cooling during the soldering operation, however, introduces significant departures from the predictions from phase diagrams. The freezing range given by the pasty gap is essentially controlled by the deviation of the alloy composition from that of
1.5 Soldering Alloys
rwrnmS
i»*H»
Figure 1-15. (a) Primary lead dendrites in a matrix of lead-tin eutectic in an alloy containing 50% Sn-50% Pb, (b) eutectic structure in a near-eutectic lead-tin alloy, and (c, d) primary tin dendrites in a matrix of lead-tin eutectic in an alloy containing 69.2% Sn-30.8% Pb.
the eutectic (61.9% Sn, 38.1% Pb). The microstructures of cast alloys consist of primary dendrites of tin-rich or lead-rich solid solutions in a matrix of lamellar or globular eutectic (Fig. 1-15 a-d). An extended freezing range allows the primary dendrites to grow and reduces the extent of eutectic solidification. The plasticity of the solidifying solders in the freezing range is influenced not only by the liquid phases present but also by the size and the shape of the primary dendrites. A narrow freezing range is normally associated with finer primary dendrites and a larger volume fraction of the eutectic liquid, which gives rise to better plasticity of the solidifying solders. The common lead-tin solders have an upper limit to their service temperature be-
cause the tensile strength drops down significantly with temperature (Fig. 1-16). Soldering alloys designed for high temperature service conditions contain a very high percentage of tin (95 % or above) and an antimony or silver addition. Table 1-4
60
100
140
180
220
TEMPERATURE , °C
Figure 1-16. Tensile strength of tin-lead alloys as a function of temperature (Mackay, 1977).
48
1 Tin
Table 1-4. Nominal chemical composition and properties of some tin solder alloys. Solder
Solder composition
Elec- Spe- Application trical cific area resis- gravtance ity (MPa) (kg/mm2) (Hfl/cm)
Melting point (°C) Tensile Shear Brinell strength strength hardness Solidus Liquidus (MPa)
1
63Sn-37Pb
—
183
54.1
38.0
17.0
14.99
8.46
2
5OSn-5OPb
183
214
43.6
36.6
14.0
15.82
8.90
3
40Sn-60Pb
183
238
38.0
33.7
12.0
17.07
9.28
12.5Sn-25Pb-50Bi-12.5Cd 70 50Sn-47Pb-3Sb 185 95Sn-5Sb 233
74 204 240
32.0 59.1 41.5
_ 48.2 42.2
25.0 15.6 13.3
55.61 17.96 _
9.50 8.75 7.20
13.7
13.7
-
— -
— -
4 5 6 7
95Sn-5Ag
221
245
56.2
-
8 9
10Sn-90Pb 5Sn-l.5Ag-93.5Pb
267 296
301 301
37.0 39.0
— -
lists the compositions of different common solders and their strengths and general usage. Alloys like 95% S n - 5 % Ag and 95% S n - 5 % Sb not only retain their strength to a temperature of 200 °C, they also possess wetting properties similar to those of near eutectic alloys. For cryogenic applications where lowtemperature strength and ductility are important factors, high lead-low tin alloys such as 90% Pb-10% Sn or 93.5% P b 5% Sn-1.5% Ag are selected. Though they exhibit poor wetting properties and inferior strength, they do not undergo the severe loss in ductility and impact strength that occurs at 173 K for body-centred tetragonal tin and tin-rich solders. The roles that different impurity elements play in influencing the wettability and other properties (Mackay, 1977) are summarised as follows: • Aluminium present at a level of 0.001 % in solder material causes oxidation and the aluminium oxide film at the
— -
Electronics and instruments Sheet-metal work General engineering
Creep resistance Creep resistance Cryogenics Cryogenics
liquid solder/flux interface can result in dewetting, particularly while soldering steels. • Antimony is present in a number of commercial solders in the range of 0 . 1 0.5 %, and in special creep resistant solders up to a level of 5%. The presence of a small percentage of antimony (0.1-0.5%) in lead-rich lead-tin solders produces a dewetting tendency on copper and brass but not on steel surfaces. An intentional addition of antimony of up to about 5 % gives rise to the formation of SbSn intermetallic particles which are responsible for imparting creep resistance to the solder material. • Arsenic, if present at a level of 0.005%, can cause de-wetting particularly on brass, supposedly due to the formation of As-Zn intermetallic compounds. • Bismuth additions of 0.1-0.25% can enhance the rate of spread of eutectic tinlead solder. When present in excess of 0.5%, a discoloration of the surface of the solder occurs. • Cadmium contamination of solder baths, which often occurs with cadmium-
1.6 Bearing Alloys
plated components, can cause a visible dulling of the solder surface due to the presence of an oxide film. The presence of cadmium decreases the wetting force and increases the wetting time quite appreciably. The increased tenacity of the surface oxide due to cadmium can give rise to soldering defects. • Copper additions in solder alloys up to about 0.5% have a small effect on the wetting properties, but solders containing more than 0.25 % copper show a gritty appearance on copper test pieces due to the presence of discrete particles of a coppertin intermetallic compound. • Phosphorus, when present at a level exceeding 0.01 %, produces de-wetting on copper and mild steel, while sulphur additions to the same level do not cause any de-wetting. The presence of SnS and PbS particles, however, gives a gritty appearance to the solder surface, and on this basis the maximum sulphur level in solders is fixed at 0.0015%. • Zinc, being a strong oxide former, gives rise to a deterioration of the surface quality of the solder when present above 0.003%. The combined effects of different impurities make it necessary to impose more stringent limits of these elements. For example, in the mass soldering of printed circuit boards, a combined total of aluminium, cadmium, and zinc is set at a maximum limit of 0.002% for antimony-free solders.
1.6 Bearing Alloys Tin-based bearing alloys can be grouped into two classes - white metals, which have to be used as a lining on a stronger shell, and alloys used as solid bearings without backings. The former class commonly
49
known as babbitts, named after Isaac Babbitt who was granted a patent in 1839, includes tin-rich (over 80% tin) and highlead alloys containing 70% or more lead and less than 12% tin. Among the solid bearing alloys, bronzes including leaded bronze and aluminium-tin ( ~ 6 % Sn) alloys are most widely used. Babbitt, while describing his invention, stated that the "boxes" which would now be called "shells or backings", can be made of any kind of metal which has sufficient strength and is capable of being tinned. The composition of the white metal quoted in the patent is also close to one of the babbitt alloys extensively used currently, 1.6.1 Property Requirements of a Bearing Alloy
In a lubricated journal bearing, the hydrodynamic pressure that builds up in the lubricant film between the moving surfaces of the bearing and the journal makes the journal "float" from the bearing. Under ideal conditions, a continuous lubricant film would completely separate the bearing surfaces and prevent metallic contact. Departures from ideal conditions resulting from misalignments, irregularities on bearing surfaces, and suspended particles in the lubricant cause local or general breakdown of the lubricant film. Even if a bearing is maintained in perfect alignment and lubricated with clean oil, breakdown of the oil film occurs on starting or stopping, i.e., when the velocity is too low for sufficient hydrodynamic pressure to be maintained. For a bearing alloy to function under these conditions it should possess the properties of conformability and embeddability. Conformability refers to the ability of a material to adapt to stresses from elastic or plastic deformations without permanently affecting the sliding behaviour, whereas
50
1 Tin
the ability of a material to embed hard particles is referred to as embeddability. A large difference in the macro-hardness of the journal and the bearing material is essential to avoid cutting and scoring of the opposing surfaces by the projections present on them. With surfaces of very different hardness values, damage is largely confined to the softer surface, which is continually smoothened by plastic flow. Protrusions in the harder surface, though not affected by the fine particles of the softer phase, are gradually lapped by hard particles embedded in the softer phase. Bearing metals based on tin do not in themselves provide a low coefficient of friction in the absence of a lubricating oil film. In fact, it is the chemical adhesion of the oil film that makes tin alloys perform so well in bearing applications. Tin-rich alloys do not generally possess good strength, and therefore it is essential to provide a stronger backing material such as bronze, steel, or cast iron. The support that the backing material provides helps reduce the flexing of the white-metal skin and consequently the possibility of fatigue failure. The load which can be supported by a bearing is limited by the following requirements: i) The bearing must operate within the range of stable lubrication under normal operating conditions, ii) The equilibrium temperature attained by the bearing must not exceed a specified value. hi) The bearing alloy must not fail by general deformation or by fatigue. The conditions for stable lubrication can be met when the value of the parameter (viscosity x speed/load) remains below a certain limit. The permissible operating temperature of a bearing is essentially controlled by the viscosity of the lubricant,
which strongly depends on temperature which in turn is a function of the operating speed. Apart from properties like sufficient strength to resist deformation by both steady and fluctuating loads, sufficient hardness to combat wear, good conformability, and embeddability, bearing alloys are required to have adequate resistance to corrosion by the lubricant and the ability to retain their strength and hardness at the operating temperature. Whitemetal bearings, which are primarily used in the as-cast form, must also bond with steel or other backing materials, and should not have a larger contraction on cooling. Some of the property requirements are of a conflicting nature. High strength to support the load and high hardness to render wear resistance do certainly not go with the requirement of good conformability. These conflicting requirements are met by a two-phase microstructure for the bearing alloy in which hard particles are distributed in a soft matrix phase. The hard particles are distributed throughout the material in such a way that they support the load while the soft matrix is worn down to a level slightly below that of the hard particles and thus provides channels for the supply of lubricant. 1.6.2 White-Metal Bearings White metals are subdivided into three major categories, i.e., high-tin alloys, highlead alloys, and intermediate alloys; the major alloying additions in these being antimony and copper. The micro-constituent phases present in tin-rich white metals can be found in the tin-rich corner of the pseudo-ternary phase diagram of tin-antimony-copper (Fig. 1-17) (Barry and Thwaites, 1983). "Matrix" refers to the solid solution of antimony in tin in which the
1.6 Bearing Alloys
51
PRIMARY a ,SbSn SECONDARY
MATRIX
+ Cu6Sn5PRIMARY ft SbSn SECONDARY t
m
I 2 COPPER (%)
Figure 1-17. The tin-rich corner of the pseudoternary Sn-Sb-Cu phase diagram relevant to whitemetal bearing alloys (Barry and Thwaites, 1983).
Cu 6 Sn 5 compound is finely dispersed, and which forms in region I of the phase diagram. The microstructure of the matrix can be designated as pseudo-eutectic. In the regions marked II, III, IV, and V coarse particles of the intermetallic compounds SbSn and Cu 6 Sn 5 are distributed. Depending on the composition, the nature and the volume fraction of the intermetallic precipitates vary. For example, alloys in region V contain only Cu 6 Sn 5 particles (Fig. 1-18) in the matrix, whereas SbSn particles are present along with Cu 6 Sn 5 in alloys in region IV (Fig. 1-19). The particles of SbSn can be easily recognised by their geometric shape, which results from the strongly anistropic surface energies associated with this phase. The presence of these hard intermetallic particles, which stand proud in the soft matrix, makes this material suitable for bearing applications. Commercial high-tin, white-metal bearing alloys have compositions that lie within the phase fields marked IV and V. The selection of compositions is based on the data on hardness, tensile and compressive strengths, ductility, and fatigue resistance of tin-antimony-copper alloys as functions of the
Figure 1-18. C'u(>Sn^ dondriles distributed in i l l " pseudo-eutectic m;tln\ in the alloy eoi
92% Sn, 5% Sb and .V\,
CIL
Figure 1-19. (a) Coarse cuboidal particles of the SbSn phase and Cu6Sn5 particles distributed in the pseudoeutectic matrix in the alloy containing 85% Sn, 12% Sb and 3% Cu. (b) Grey regions in the matrix represent the cored solid solution of antimony in tin.
copper and the tin content. Figure 1-20, which summarises the mechanical properties of these alloys (Hedges et al., 1960), indicates that the rise in strength and hard-
52
1 Tin
7 Cu
21.7
^--'*~
- — ,3.5 Cu 'ICu
•
o 18.6
Q_
^ ^ ^
/
-^OCu ^^<^
^ ^ ^
l"l5.5 0 % Cu
2 12.4 4 % Cu 3 % Cu 2 % Cu I % Cu
fe
UJ Z>
9.3 /
/
s
I
I
P6.2 /
3.1 /
10 0
n 0
2
4
6 6 10 12 14 16 18 ANTIMONY PER CENT, %
4
6 8 10 12 14 16 18 ANTIMONY PER CENT %
,
1
2
.
1
,
1
,
1
.
1
,
1
.
1 .
4 6 8 10 12 14 ANTIMONY PER CENT, %
2
3
4
5
6
16
7
COPPER PER CENT, %
420°C
at 100 C
4
6 8 10 12 14 16 18 ANTIMONY PER CENT, %
Figure 1-20. (a) Ultimate tensile strength, (b) elongation, (c) hardness, and (d, e) fatigue strength of Sn-Sb-Cu alloys.
1.6 Bearing Alloys
ness is quite sharp when 1 % copper is added to pure tin and when the antimony level is raised to 6-8%. The former effect can be attributed to the introduction of the primary Cu 6 Sn 5 intermetallic into the microstructure while the latter corresponds to the presence of the primary Cu 6 Sn 5 and the secondary SbSn intermetallic phases. The combination of these two intermetallic phases in the form of primary Cu 6 Sn 5 needles providing a network in which SbSn cubes are entangled is basically responsible for the desirable properites of these alloys. The fatigue strength of binary tin alloys exhibits a sharp rise with an addition of about 1% copper and in ternary alloys containing antimony the fatigue strength attains a saturation value on the addition of 1 % copper (Fig. l-20e). In contrast, the fatigue strength can be enhanced with an increasing addition of antimony up to a level of about 8%. The presence of the primary intermetallic phase Cu 6 Sn 5 in the former and SbSn in the latter does not therefore contribute much towards the improvement of the fatigue strength. Antimony in solid solution with tin and the pseudo-eutectic structure consisting of alpha tin and Cu 6 Sn 5 are both responsible for improving the fatigue strength. Small additions of other alloying elements such as bismuth, tellurium, and cadmium increase the fatigue strength of white metals. The lead-base bearing alloys are based upon the lead-antimony-tin system having a microstructure containing hard particles distributed in a soft matrix. In one class of alloys containing 12-18% antimony and 5% tin, an antimony-rich solid solution acts as the hard phase in the soft matrix of the ternary eutectic (Fig. 1-21). The other class of alloys which contain 10-12% tin and about the same level of antimony consist of hard cuboid particles of SbSn in the pseudo-binary eutectic of lead and SbSn.
53
Figure 1-21. Sb-rich particles distributed in the matrix of the ternary eutectic of Sn-Sb-Pb in the alloy containing 79% Pb, 15.5% Sb and 5.5% Sn.
The latter class of alloys with a higher tin content have better compressive strength, fatigue strength, and ductility. This is in line with the fact that lead-base alloys are a cheaper substitute for tin-base alloys, and, in general, the required mechanical properties of lead-base alloys are inferior to those of tin-base alloys. Other major shortcomings of lead-base alloys include their tendency for segregation, in which the low density, primary phase particles rise up from the heavier, high-lead matrix, their poor thermal conductivities, and their softening tendency at elevated temperatures. The intermediate alloys containing 2050 % tin are inferior to both high-tin and high-lead alloys as far as elevated temperature strength is concerned. However, the wear resistance of these alloys is certainly better than that of their lead-rich counterparts. The main reason for the choice of intermediate alloys for diverse applications such as bearings in rolling mills, electric motors, and crankshafts is their lower cost and satisfactory room temperature properties. Since the matrix in these alloys is a peritectic complex that melts at about 180°C, their use is severely restricted at elevated temperatures.
54
1 Tin
White metal alloys are almost always bonded to a strong backing for use as bearings. In order to establish a metallurgical bond between the backing shell and the bearing lining, shells which are made of bronze, gun metal, steel, or cast iron are pre-tinned by hot dipping in molten tin. Intermetallic compounds form at the interface between tin and the base metal. The thickness of the brittle intermetallic compound layer is kept to a minimum by restricting the interaction time and the temperature during the tinning operation. White metal is cast in the pre-tinned shell, ensuring the following: i) The tin layer on the pre-tinned backing should remain completely molten when the white metal is poured, ii) The white metal should solidify from the backing interface so that shrinkage cavities do not form on this interface, iii) The rate of cooling is controlled, depending on the size of the shell and the composition of the white metal. The metallurgical bond between the pretinned backing and the white-metal lining is established through the formation of Cu 6 Sn 5 at the bond interface. The presence of the FeSn2 layer on the steel backs and of the Cu 6 Sn 5 layer on the bronze backs is responsible for nucleating Cu 6 Sn 5 from the molten white metal. Rapid cooling, however, is necessary to prevent the formation of a continuous layer of the brittle Cu 6 Sn 5 phase at the interface. After the solidification is complete, a slow cooling down of the shell to room temperature is commonly practiced to allow the relief of thermal stresses in the white metal. 1.6.3 Aluminium-Tin Bearing Alloys
Aluminium-tin alloys for bearing applications can be grouped into two classes, namely, low-tin alloys containing 6-7%
tin and small amounts of copper, nickel, silicon, or magnesium and high-tin alloys which contain 20-40% tin and about 1 % copper. The low-tin aluminium bearing alloys are known for their high fatigue strength and seizure resistance, but their embeddability is not very satisfactory. They tend to give high wear when used against soft shafts unless a very clean lubricant is used. Since the thermal expansion coefficient of aluminium-rich alloys is about twice that of steel, special design features must be incorporated to prevent the highaluminium bearings from becoming loose in their steel housings. Bearing alloys containing 20-40% tin show a microstructure in which tin is distributed in the form of a continuous intergranular network in the as-cast condition. These alloys are generally used in the coldrolled and annealed condition. The intergranular layer of tin is broken up during the cold-working operation and the aluminium matrix is recrystallised on subsequent annealing. These bearing alloys are usually bonded to steel backings by rolling onto aluminised steel. 1.6.4 Bronze Bearings
The high strength level of phosphor bronze, a typical composition being about 10% tin and 0.5% phosphorus, makes it suitable for applications in which heavy loads and high temperatures are involved. An addition of lead to these alloys (leaded bronzes) improves conformability at the expense of strength and wear resistance. Sintered bronze bearings made by an appropriate powder metallurgy route can retain connected porosity to an extent necessary for maintaining lubricant in these pores. These are widely used in small machinery and do not require regular oiling.
1.7 Other Alloys
Sintered bronze bearings are impregnated with Teflon or other plastic material for making oil-free bearings. 1.6.5 Underwater Bearings Bearings that encounter the service environment of steam or aerated seawater and are galvanically coupled with a large mass of steel are commonly made of an admiralty underwater bearing alloy having a nominal composition of 68.5% tin, 30% zinc, and 1.5% copper. The microstructure of this alloy consists of a distribution of CuZn 3 and the zinc-rich terminal solid solution in a matrix of tin-zinc eutectic. This alloy, which remains cathodic with respect to mild steel, corrodes preferentially compared to the shaft.
1.7 Other Alloys 1.7.1 Tin-Rich Alloys 1.7.1.1 Die-Casting Alloys The low melting temperatures and very high fluidity of tin alloys have made them suitable for die casting for the production of sound castings of intricate designs. Dimensional accuracy to a level of ± 0.02 % can be achieved because of the low working temperature and low shrinkage. Apart from the ease of manufacture into thin sections and intricate shapes, these alloys have good bearing properties and wear resistance in unlubricated positions. A wide range of compositions of tin alloys containing the alloying elements antimony, copper, and lead are used in die casting; the choice of alloy being decided primarily on the basis of property requirements.
55
1.7.1.2 Pewters Pewters, a range of alloys containing about 90% tin and alloying additions of antimony, copper, and lead have been in use since Roman times for utensils, decorative items, and costume jewellery. Modern pewters which are free of lead do not lose their bright finish with age. Traditionally, pewter articles were cast and the more expensive heavy articles are still produced by casting. Deep drawing and spinning techniques are increasingly used for the manufacture of pewter articles from rolled sheets. The spinning of pewters is attractive because of the fact that these alloys do not exhibit work hardening. 1.7.1.3 Fusible Alloys The range of alloys based on eutectics of low melting metals, commonly known as fusible alloys, correspond to alloys with fairly regularly-spaced melting temperatures from 20 °C to 227 °C. Depending on the nature of the applications, namely, fire alarms and fuses, seals, holding of workpieces or tools for machining, temperature indicators, fabrication of plastics and coating moulds, suitable alloys can be chosen which melt or solidify at the specified temperature level. In addition to the binary and multicomponent eutectics based on tin, bismuth, lead, cadmium, indium, and gallium, some of the eutectic alloys are also included in the range of fusible alloys. Alloys containing a high proportion of bismuth are special as they expand on freezing. Control over the volume change during freezing is therefore effected by the addition of bismuth to an appropriate level. 1.7.2 Tin as a Minor Additive Tin is added to a number of important alloys as a minor but vital constituent. The
56
1 Tin
importance of tin in some of these alloys is highlighted in this section. 1.7.2.1 Tin Bronzes Primary alloys of copper and tin are possibly the oldest alloys known to mankind. It has been suggested that the first production of tin bronzes may have involved the direct smelting of copper ores containing naturally occurring tinstone. Apart from the use of bronze for tools, weapons, and other utilitarian purposes, it has remained an alloy for the sculptor for many centuries. The physical metallurgy of tin bronzes can be briefly discussed using the copperrich portion of the binary copper-tin phase diagram (Fig. 1-22). The copper-rich terminal solid solution (f.c.c. structure), shows a maximum solubility of 16.2% tin at 520 °C, the solubility sharply decreasing
10 20 30
40
with temperature. However, the retention of about 13% tin in the supersaturated a-solution is not very difficult because of the sluggish kinetics of precipitation of the intermediate phases from a. The (3-phase (b.c.c. structure) forms as a result of a peritectic reaction at 798 °C and decomposes in a eutectoid reaction to (a + y) at 586 °C. The y-phase in turn undergoes a eutectoid decomposition into (a+ 5), the 5-phase being a typical electron compound having the stoichiometry of Cu 31 Sn 8 and the structure of y-brass. The 8-phase, which is not a primary solidification product, is frequently present in tin bronzes as a component of the eutectoid decomposition product. The phase diagram also indicates the eutectoid decomposition in which the 5-phase decomposes into a mixture of oc and 8 (h.c.p. structure). This reaction, however, is so sluggish that in commercial tin-bronzes the 8-phase is
WEIGHT PERCENT TIN 50 60 70 80
100
1100
UJ
a
Ior
UJ
a. UJ
200 100
0 Cu
10
20
30
40 50 60 70 ATOMIC PERCENT TIN
80
Figure 1-22. The equilibrium copper-tin phase diagram (Massalski, 1986).
90
100 Sn
1.7 Other Alloys
not generally present. The solvus line below the eutectoid temperature of 350 °C shows a marked decrease in the solubility of tin in copper. However, the sluggish separation of the s-phase from a is responsible for the retention of a sufficiently high level of tin in solid solution in industrial alloys. Slow cooled, cast tin bronzes containing below 7 % tin generally show only a single phase oc-structure, whereas most castings containing over 7% tin contain the S-phase distributed within the gaps between the oc-dendrites. Ternary and quaternary alloys based on bronzes contain phosphorus, zinc, or lead. The phosphor bronzes are characterised by high strength, toughness, high corrosion resistance, low coefficient of friction, and resistance against season cracking. Phosphor bronzes are used extensively for diaphragms, bellows, lock washers, cotter pins, springs, etc. Zinc is sometimes used to replace part of the tin. The result is an improvement in the casting properties and toughness. Lead is often added to tin bronzes to improve machinability and wear resistance. High-lead tin bronzes may contain up to a maximum of 25 % lead. The leaded alloys are used for bushings and bearings under moderate or high loads.
gam from an alloy (silver-tin or silvertin-copper) with equal amounts of mercury using a mortar and pestle and filling the tooth cavity under a pressure sufficient to squeeze out any excess mercury. A long lasting, hard restoration can be achieved which shows the required dimensional stability and mechanical and corrosion properties. The procedure for dental filling was developed mostly through meticulous experimentation at a time when neither the knowledge of the silver-tin phase diagram was available nor the phase analysis through microscopic and diffraction techniques was known. The metallurgy of the process, which was understood only after the availability of modern characterisation tools, can be briefly described as follows: The 70-30 silver-tin alloy corresponds to the y-phase (ordered orthorhombic), as shown in Fig. 1-23. The y-phase reacts with the mercury to produce a silver compound (y1) and a tin-mercury compound (y2)- These two reaction products bond partially-reacted particles of the y-phase. It is this process which causes the amalgam to harden.
1.7.2.2 Tin in Dental Amalgam The restoration of decayed teeth with amalgams prepared from silver-tin alloys has been practiced for over a hundred years. In spite of the concern over possible mercury toxicity and of the availability of composite resin dental filling material, present-day dentists realise that there is no alternative to dental amalgam which has been proven to give durable restorations at reasonable costs. The procedure of filling a tooth cavity involves the preparation of a dental amal-
Sn
10 20
30
40
50
60
70
80
90
Ag
WEIGHT % A g
Figure 1-23. The equilibrium tin-silver phase diagram (Hansen and Anderko, 1958).
58
1 Tin
The reaction of annealed filings of a y-phase silver-tin alloy with an equal mass of mercury continues over a period of five or six hours and the process is accompanied by a slight volume contraction. On the other hand, the silver-rich ^-phase (disordered hexagonal close-packed) reacts rapidly with mercury and undergoes an expansion of 0.5-1.0% as it hardens. It is for this reason that an appropriate composition of the alloy is chosen. Close control of the dimensional change during hardening of a mixture of the y- and ^-phases of silver-tin alloys is therefore possible. The durability of the restoration is strongly dependent on the mechanical and electrochemical properties of the hardened amalgam. The y2-phase (Sn7Hg) has been shown to have an adverse influence on the durability (particularly the corrosion resistance) of amalgam restorations. The addition of small amounts of copper to the amalgam alloy is known to suppress the formation of the y2-phase by producing Cu3Sn (the y-phase of the copper-tin system), which when present as finely dispersed precipitates contributes significantly to the strengthening of the amalgam. The development of ternary Ag-Sn-Cu amalgam alloys with a composition of typically 60 % Ag-27 % Sn-13 % Cu represents a major advance in the durability of amalgam restorations. The use of rapidly solidified atomised particles of amalgam alloy is another advance made in recent times. Studies on the mechanism of the reaction of the ^-phase and the y-phase with mercury has revealed that the mercury attack is, in general, uniform on the tj-phase alloys, while for the y-phase preferential attack occurs at grain boundaries, slip bands, and deformation twins, i.e., at regions of high dislocation density and low degree of order in an otherwise ordered structure. The annealing of filings of the
amalgam alloys before reacting with mercury is essential in order to reduce these sites to slow down the preferential attack, thereby allowing the dentist ample time to condense the pasty amalgam into the tooth cavity. 1.7.2.3 Superconducting Nb3Sn The intermetallic compound Nb 3 Sn exhibits the A15 structure (Cr3Si type), which is cubic with the space group Pm3n. The A15 compounds are associated with A 3 B stoichiometry, where A is one of the transition elements from the titanium, vanadium, and chromium columns and B belongs to the later columns or nontransition elements. In the unit cell 6A atoms are located at positions with a point symmetry of 42m while 2B atoms are located at the cell corners and the body-centre position. The short A-A distance along the three mutually orthogonal chains on the unit cell faces is the characteristic feature of this structure. Many of the A15 compounds exhibit attractive superconducting properties (see Volume 3 A of this Series, Sec. 4.4.6). Amongst them Nb 3 Sn occupies an important position, as can be seen from a comparison of the superconducting properties, i.e., critical temperature, Tc9 and upper critical field, Hc2 (Table 1-5).
Table 1-5. Comparison of superconducting properties of important A15 compounds. Compound
Tc (K)
Hc2 at 0 K (T)
V3Ga V3Si Nb3Sn Nb3Al Nb 3 Ga Nb 3 Ge
14.8 16.9 18.0 18.7 20.2 22.5
34.9 34.0 29.6 32.7 34.1 37.1
1.7 Other Alloys
As for other type II superconductors, the achievement of a high critical current density Jc depends crucially on the microstructure of the A15 phase. Efficient flux pinning, which promotes the enhancement of Jc, requires a distribution of the inhomogeneities on a scale comparable with the periodicity of the flux line lattice. It has been established that the Jc of Nb 3 Sn in high magnetic fields is inversely proportional to its average grain diameter until the grains become less than SOSO nm. In order to maximise the grain boundary area for more effective flux pinning, special methods like doping with ZrO 2 and synthesis of the compound at low, solid state reaction temperatures (550-750 °C) are employed and finegrained (<100nm) Nb 3 Sn is produced with a Jc of - 1 0 6 A/cm2 at 5 T. Of the many known superconducting A15 compounds, Nb 3 Sn and V 3 Ga have the distinction of being produced on a commercial scale. This is due to the feasibility of synthesising these compounds by several novel techniques, i.e., liquid-solute diffusion, chemical vapour deposition, physical vapour deposition and sputter deposition, and solid state diffusion. Liquid-solute diffusion is one of the initial processes adopted for the commercial production of Nb 3 Sn tapes for high field magnets. The process consists of the introduction of thin niobium tape (~25 jam thickness) into a molten tin bath. The tincoated tape is then heated at 930-1100 °C in a vacuum or in an inert atmosphere to form Nb 3 Sn layers on both sides of the tape. The most important parameter in this process is the amount of tin left on the niobium after the niobium tape has passed through the tin bath. In one variation of the process, the thickness of the tin layer is so controlled that after the reaction treatment the entire tin layer is consumed in
59
forming the dense Nb 3 Sn phase. The other variation involves the deposition of excess tin on the tape, which allows the growth of thicker Nb 3 Sn layers and the penetration of excess liquid tin into the Nb 3 Sn grain boundaries. In the second process further surface treatments are necessary to reduce ac losses. Copper can be soldered directly to the reacted ribbon for electrical stabilisation, while the entire tape can be encased in stainless steel for protection and mechanical strength. The chemical vapour deposition technique consists of the production of gaseous chlorides by passing chlorine over the metallic elements at 800-900 °C, and hydrogen reduction of the mixture of chlorides of constituent elements onto a heated substrate of niobium, stainless steel, or Hastelloy. The superconducting properties of these compounds are very sensitive to factors like deposition temperature and flow rates. In general, various impurity gases such as O 2 , CO, and CH 4 are added to improve the critical current density. Co-deposition of constituent elements by electron beam and sputtering techniques provides a method of producing Nb 3 Sn deposits. Control of the substrate temperature is vital in controlling the grain size of the deposit. Fine layering of Nb 3 Sn with nonsuperconducting materials and thereby enhancing the Jc is possible through this technique. Solid state diffusion, popularly known as the bronze process, has been developed for the production of Nb 3 Sn multifilamentary conductors. The process involves the insertion of niobium rods in a tin-bronze billet. These composite billets are reduced to small diameter wires by extrusion and wire drawing. The wires are rebundled and again packed in a copper or bronze tube for further drawing into a multifilamentary composite. Depending upon the
60
1 Tin
usage, the composite wire is given a final mechanical twisting so that the filaments take up a helical path which assists in achieving electrical stability in changing magnetic fields. Formation of the superconducting Nb 3 Sn phase is induced by heat treatment at 700 °C for several hours. The overall solid state process involves selective diffusion of tin from the bronze matrix to the surfaces of niobium filaments and the formation of the A15 compound. The heat treatment operation can precede or follow the winding of the conductor into a solenoid shape. The formation of Nb 3 Sn in a solid state reaction between tin bronze and niobium (Fig. l-24a-d) is possible due to the fact that a ternary equilibrium is established between the Nb 3 Sn phase, the niobium solid solution, and the copper solid solution at temperatures close to 700 °C. An isothermal section of the Nb-Sn-Cu phase diagram at 700 °C (Fig. l-24e) shows the diffusion path (dashed line) between the Nb-Sn solid solution and the Cu-Sn solid solution (ss), which passes through the Nb 3 Sn phase field. 1.7.2.4 Tin in Titanium Alloys Tin is soluble in both the |3- (b.c.c.) and the a- (h.c.p.) titanium-rich terminal solid solutions to fairly large extents, the maximum solubilities being about 36.7% and 27%, respectively. The (3 to oc transformation temperature of titanium is not affected significantly by the addition of tin. Solid solution strengthening of the oc-phase by tin is utilised in a number of commercial a- and (<x + (3) alloys of titanium. If the tin content exceeds the solubility limit in the ot-phase, an ordered intermetallic phase Ti3Sn (D0 19 structure) forms. Hardening by this intermetallic phase, however, is not practised in commercial titanium alloys.
1.7.2.5 Tin in Zirconium Alloys The most commonly used zirconium alloys for structural materials in nuclear reactors contain tin as a vital alloying addition. These alloys, with the generic name of zircaloys, utilise the influence of tin in improving both the mechanical and the corrosion properties of zirconium alloys without seriously affecting their low neutron absorption cross sections (see Volume 10 B of this Series, Chap. 7). The corrosion resistance of a-zirconium (h.c.p. structure), which is significantly improved by an addition of tin, is believed to be due to its ability to compensate for the adverse effect of nitrogen. The role of tin as a solid solution strengthener and a precipitation hardener in zirconium-base alloys has recently been examined (Wadekar et al., 1991). The mechanical properties of binary Zr-Sn alloys containing 0.5-8% Sn in the betaquenched and tempered conditions are shown in Fig. 1-25. The 0.2% yield stress of solution-treated alloys shows a linear relationship with C 2/3 , where C represents the atomic concentration of tin. This suggests operation of the Mott-Nabarro model of solution hardening in which the clustering of solute atoms is envisaged. Transmission electron microscopy of solutiontreated dilute alloys (C<0.04) indeed shows the presence of strain contrasts associated with the solute clusters (Fig. 1-26 a). It has also been shown that an addition of tin influences both the athermal and the thermal components of the flow stress. Precipiation of the Zr3Sn (A15 structure) phase, which occurs in dilute alloys (C<0.04) on aging and even in solution-treated alloys of higher tin contents, is responsible for the hardening. Partial coherence between the precipitate phase and the matrix is maintained with the basal
1.7 Other Alloys
61
(e)
4-AI5
Figure 1-24. (a) Cross section of a multifilamentary Nb3Sn composite superconducting wire, (b) The formation of Nb3Sn at the interface between niobium filaments and the bronze matrix due to the solid state reaction, (c) Secondary electron image showing the interface reaction between a bronze filament and the niobium matrix, (d) X-ray image of the same area shown in (c) formed by superimposition of images due to NbK a and SnKa. The brighter periphery corresponds to the Nb3Sn layer formed at the interface (courtesy of M. C. Saxena, unpublished work), (e) Niobium-rich corner of the isothermal section at 700 °C of the Nb-Sn-Cu ternary phase diagram (Luhman and Dew Hughes, 1979).
plane of the h.c.p. alpha-zirconium remaining parallel to the (111) plane of Zr3Sn (Fig. 1-26 b).
diagram. About 17.7% tin is soluble in oc-iron at 900 °C and this falls to about 10% at 500 °C. Tin causes serious embrittlement of steels, and both cold and hot shortness, and is therefore considered a very detrimental impurity. The tin content in finished steels is generally specified as below 0.05 %. The control of tin is particu-
Nb s s + Cu
AI5
1.7.2.6 Tin in Steels and Cast Irons Tin, being an a- (b.c.c.) stabiliser, gives rise to a y- (f.c.c.) loop in the iron-tin phase
62
1 Tin
700
20
Q Q+T o-U.T.S.- • A- Y.S.-A v-Elong.-*
JQ~
/
K7
600 -
u
/
Ui 500 (/) UJ
\(J)
XT
A/,,
/ /
400
- 10 I
A -
-5
1
1
2
1
1
1
4 6 TIN (Wt%)
8
Figure 1-25. Yield stress (Y.S.), ultimate tensile stress (U.T.S.), and percentage elongation of zirconium-tin alloys as a function of the tin content for two types of microstrutures, namely, beta-quenched (p-Q) and beta-quenched and tempered (Q + T) (Wadekar et al., 1991).
larly important in secondary steel making processes in which tin plate may sometimes be used as a constituent of the scrap feed. Limiting the content of tin plate would not only improve the quality of the steel, but tin plate scrap could also be used for the recovery of tin and the residue scrap used as a charge in steel making. Tin is known to be a strong pearlite stabiliser in both grey and nodular cast irons. Small additions of tin improve the wear resistance of cast iron by ensuring a completely pearlitic matrix without forming massive cementite which causes machining difficulties. The stability of cementite with respect to its decomposition into graphite and ferrite improves with small additions of tin. This property is utilised in cast irons intended for use at elevated temperatures.
1.8 Tin Compounds 1.8.1 General
0 4 ,iim y.
.„..
.
.
•
_...f
Figure 1-26. Contrasts seen in TEM micrographs from solute clusters and Zr3Sn precipitates in the Zr-8% Sn alloy in conditions (a) beta-quenched, and (b) beta-quenched and tempered at 650 °C for 30 min (Wadekar et al., 1991).
A large number of inorganic and organic compounds of tin are used for a variety of applications. Cumulatively, these compounds form the third largest demand sector for tin after tin solder and tin plate. The use of inorganic compounds at present is more than that of organic compounds, but the rate of growth in the applications of organic compounds is quite high. The application areas of tin compounds include their use as catalysts, stabiliser for plastics and polymers, opacifiers in enamels and ceramic glazes, additives in food and oral hygiene products, additives in paints, biocidal agents, wood preservatives, additives in certain chemicals for treating textiles, as electroconductive coatings, as sensors for analysing certain gases, and as fire retardants and smoke suppressants.
1.8 Tin Compounds
1.8.2 Commonly Used Inorganic and Organic Compounds
Some of the commonly used inorganic compounds are stannous and stannic chlorides, stannous and stannic oxides, stannous fluoride, sodium and potassium stannates, stannous fluoborate, and stannous sulphate, but over 80 % of the consumption of tin compounds is in the form of tin chlorides and oxides. The physical and thermodynamic properties and the most common preparative process for some of the commonly used inorganic compounds are given in Table 1-6. Organic tin compounds can be broadly classified into two groups, one group comprises compounds with organic acids in which the tin forms bonds with radicals and not directly with carbon, while in the other, known as organotin compounds, tin forms at least one tin-carbon bond. Some of the commonly used organic compounds are tin acetate (CH3COO)2Sn, stannous ethyl hexoate Sn(C 8 H 15 O 2 )2 termed octoate, stannous formate Sn(CHOO) 2 , stannous oxalate Sn(C 2 O 4 ), and stannous tartarate Sn(C 4 H 4 O 6 ). Organotin compounds are tetra-(R4Sn), tri-(R3SnX), di(R 2 SnX 2 ) and mono-(RSnX3) type compounds in which R is methyl, butyl, octyl, cyclohexyl, or phenyl and the noncarbon bonded group X is halide, oxide, hydroxide, carboxylate, etc. The manufacture of organotin compounds usually involves first the synthesis of a tetra-organotin compound by direct reaction of stannic chloride with organoaluminium, organomagnesium, or as organic compound in the presence of sodium. The tetra-compound is then further reacted with stannic chloride to obtain tri-, di- and mono-organic chloride compounds. Other compounds are then synthesised from the respective organotin chloride compound.
63
Amongst these compounds tetra- and mono-organotin compounds have very limited direct commercial applications. The tri-organotin compounds are mostly used for biocidal applications and the diorganotin for a variety of other applications. Some of the physical properties and common methods of preparation are shown in Table 1-7. The application areas of inorganic, organic, and organotin compounds are briefly described in the following sections. 1.8.3 Applications of Compounds 1.8.3.1 Inorganic Compounds
Halides Tin chlorides are available in anhydrous and hydrated forms. Anhydrous compounds, though somewhat difficult to handle, are used when moisture is undesirable in their applications. Commonly used halides are stannic chloride, stannous chloride, and stannous fluoride. Anhydrous stannic chloride is used in the manufacture of most of the organotin compounds, in catalysing some of the organic reactions such as alkylisation, cyclisation, esterification, halogenation, and polymerisation, for stabilising colour in soaps, in improving the uptake of colour in the dyeing of silk, and as an antistatic agent for synthetic fibres. It is now used extensively for applying coatings of tin oxide by decomposition of the tetrachloride at 500-600 °C on hot glass. Hydrated tin chloride (SnCl4 • 5 H 2 O) is used as an additive in formulations for fire proofing wool. Stannous chloride is used in electroplating tin and tin alloys, as a reducing agent in many of the inorganic and organic redox reactions, as a sensitising agent to provide a metallic surface prior to metallic coating on many nonconducting surfaces
64
1 Tin
Table 1-6. Physical and thermodynamic properties and preparative methods for some commonly used inorganic compounds of tin. Compound
Colour, form
Density
Boiling point (°C)
— AG298
— AH 298 Preparative method
(g/cm3)
Melting point (°C)
(kJ/mol)
(kJ/mol)
SnCl2
White, crystalline
3.95
246.7
605
207.16
197.76
SnCl4
Colourless, liquid White, crystalline
2.23
-30.2
114
431.66
471.09
Blue black, crystalline White, crystalline
6.45
256.95
285.70
SnO2
White, crystalline
6.9
519.54
580.27
SnO2xH2O
White
SnF 2
Opaque, white
4.9-5.3 219.5
491.87
493.49
Na 2 SnO 3 -3H 2 O
Colourless, crystalline Colourless, crystalline Solution
3.03
White, crystalline
4.18
SnCl 4 -5H 2 O
SnO SnO • H 2 O
K 2 SnO 3 -3H 2 O SnBF4
SnSO4
2.04
—
-1600
— 853
3.30
Stable in temperature range of 19-56 °C;
3
2
decomposes above 385 °C;
3
Heating tin in HC1 gas or by direct chlorination Direction chlorination of tin at 110-115 °C Dissolving SnCl4 in hot water and crystallizing by cooling Precipitating with alkali from SnCl2 Adding alkali solution to a stannous salt solution Oxidising tin metal with air or burning fine powder, by calcining hydrated oxide or by dechlorinating SnCl4 with steam, treating tin with nitric acid or by neutralising SnCl4 with alkali By hydrolysis of stannate compounds Dissolving SnO in HF or by reaction of tin with HF By fusing SnO 2 with NaOH or Na 2 CO 3 By fusing SnO2 with KOH or K 2 CO 3 Dissolving SnO in fluoroboric acid or by electrolysis of tin in fluoroboric acid Reacting tin with excess sulphuric acid or by direct electrolysis of tin in sulphuric acid in cells with anion exchange membranes
decomposes above 360 °C.
1.8 Tin Compounds
including the coating of plastics and the silvering of mirrors, in the coating of sensitising paper for blueprints, and as an additive to many canned food items. It catalyses many of the organic reactions such as hydrocarbon conversion, polymerisation, condensation, curing of resins and rubber, esterification, halogenation, hydrogenation, oxidation, etc. It is also used as an additive for stabilising the perfume in toilet soaps and as an anti-sludge agent in oils. Stannous fluoride is a soluble, nontoxic compound for incorporating fluoride ions in oral hygiene products such as toothpastes for preventing the demineralisation of teeth and for preparing mouthwash, chewing gum, etc. Oxides Stannous oxide is essentially used as an intermediate in the preparation of stannous compounds. A certain amount is, however, used in the preparation of goldtin and copper-tin ruby glass. Stannic oxide is a nonstoichiometric, oxygen deficient compound with an n-type superconducting behaviour. In normal oxidising conditions, the conductivity of the oxide is reduced due to the presence of chemisorbed oxygen on the surface, which withdraws electrons from the conduction band of the bulk solid. The conductivity of the oxide is therefore dependent on the temperature and the environment. Many of the recent uses of tin oxide are based on these properties. Tin oxide is used as a melting electrode in the glass making industry because of its good electrical conductivity at high temperatures and its chemical compatibility with molten glass, it also permits operation of the furnace at a somewhat lower temperature, thus reducing the loss of volatiles
65
and facilitating cleaner operation. Tin oxide electrodes have some attractive features for application in certain metallurgical industries. Since the conductivity of the oxide is dependent on the proportion of chemisorbed oxygen, any change brought about by the presence of gases such as carbon monoxide, methane, other hydrocarbons, alcohol, or ammonia would alter the amount of chemisorbed oxygen and thus affect the conductivity of the oxide. This change in conductivity can be calibrated with respect to the composition of the gaseous environment. It is therefore used for making devices for the detection and quantitative determination of many of the gases. The sensitivity and selectivity of devices can be controlled by varying the temperature and/or introducing some dopants or additives. A coating of 100 nm thick tin oxide on glass is optically transparent and electrically conducting making it suitable for use as a conducting layer on the glass surface for use as windshields in aircraft, lighting panels, display signs, fluorescent lights, or for electron beam control in cathode ray tubes. Moreover, the oxide layer enhances the strength and abrasion resistance of the glass, making it possible to use thinner glass. One of the other uses of tin oxide is as a component in ceramic dielectrics and capacitors. Since tin oxide is insoluble in many of the glaze compositions in the ceramic industry, it is used as an opacifier, as a white glaze, and as a base for carrying pigment such as chrome-tin pink, vanadium-tin yellow, and antimonytin blue-grey. Tin oxide is also used as a grinding and polishing media for marble, granite, glass, plastic lenses, and sewing needles. Hydrated stannic oxide has good potential for use in the field of ion exchange and as a catalyst with or without combination with other oxides for some of the oxidation reactions.
Table 1-7. Physical properties and common methods of preparation of some organic compounds of tin. Compound
Chemical formula
Molecular weight
Physical form
Specific gravity
Stannous acetate
Sn(CH3COO)2
236.69
Colourless, crystalline
Stannous formate
Sn(CHOO)2
208.69
Anhydrous, crystalline
Stannous oxalate
Sn(C2O4)
306.69
Stannous tartarate
Sn(C4H4O6)
266.69
White, crystalline White, crystalline
Stannous 2-ethylhexoate (Stannous octoate)
Sn(C8H15O2)2
404.69
Light yellow, viscous liquid
Tetrabutyltin
Sn(C4H9)4
347
Oily liquid 1.05
Tetraphenyltin
Sn(C6H5)4
427
White, crystalline
Melting point
Boiling point
Application areas Method of preparation
Organic compounds
Catalyst in hydro- Stannous chloride in glacial acetic acid in genation the presence of oxygen Reacting stannous Catalyst for oxide with formic acid hydrogenation of liq. fuels Catalyst for esterification reactions Dyeing and print- Reacting stannous ing of textiles chloride and potassium formate Stannous oxide with Production of 2-ethylhexoic acid polyurethane foams and room temperature vulcanisation of silicones
182
Organotin compounds
-97
1.48-1.49 224-230
145
>420
Corrosion inhiReacting SnCl4 with biting addition to excess butyl magnesium chloride in lubricating oils toluene above 100 °C, washing and purification Prepared by GrigUsed in chlorinard's reaction, i.e. nated dielectric fluids for scaveng- SnCl4 with phenyl magnesium chloride ing acids in tetrahydrofuran
Bis(tributyltin)oxide
(C4H9)3SnOSn(C4H9)3 596
Tributyltinchloride
(C4H9)3SnCl
325.5
Colourless, 1.2-1.3 liquid
Triphenyltinhydroxide
(C6H5)3SnOH
367
White, powder
120
Triphenyltinchloride
(C6H5)3SnCl
385.5
White, crystalline
105
Dimethyltindichloride
(CH3)2SnCl2
220
106
Dibutyltindilaurate
3
632
Colourless, — crystalline Liquid 1.05
Dimethyltinbis(isooctylmercapto acetate) Dibutyltinmaleate
4
491
Liquid
5
322.5
White, powder
1
Decomposes above 100°C;
2
decomposes above 280°C;
Colourless, 1.17 liquid
3
<-45
30
23
1.18
Heating a mixture of tributyltinchloride and aqueous sodium hydroxide 142 (20 Pa) Used for making By heating tetrabutylother tributyltin tin with stannic chlocompounds and as ride at 210-230 °C a rodent repellant in rubber and plastics — Fungicide From corresponding chloride by alkali hydrolysis — Antifouling paint Heating tetraphenyltin formulations, and stannic chloride preservative for marine timbers 185 -190 Glass coating Reacting methylchloprecursor ride with tin 400 (13 Pa) Stabiliser for PVC, Reacting dibutyltincatalyst for oxide with lauric acid polyurethane foams and cold curing silicone elastomers PVC stabiliser PVC stabiliser, Reacting dibutyltinsuitable for pack- oxide with maleic aging foodstuff anhydride in methanol at 60-80°C
103
(C 4 H 9 ) 2 Sn(OOCC n H 23 ) 2 ;
212 (13 Pa) Biocide, wood preservative
4
(CH 3 ) 2 Sn(SCH 2 CO 2 C 8 H 17 ) 2 ;
5
(C4H9)2Sn(C2H2O4).
O o •D
o c
13
j
68
1 Tin
Other Compounds
1.8.3.2 Organic Compounds
Sodium and potassium stannates are used mainly for electroplating tin for tin plate and for alloy coatings. Potassium stannate is preferred by many of the plants because it offers the advantage of a higher cathode current efficiency in electrodeposition compared to sodium stannate as an electrolyte. Moreover, the solubility of potassium stannate is also high at 1105 g/1 compared to 615 g/1 for sodium stannate at 15 °C. The solubility of potassium stannate increases with an increase in the temperature, whereas that of sodium stannate is reduced. Sodium stannate is also used as an additive in fire retardation formulations. Zinc stannate (ZS) and zinc hydroxystannate (ZHS) are very effective fire retardants for synthetic polymers and smoke suppressants for synthetic polymers and paints. These compounds are used as additives in a series of hydrogenated organic compounds to impart useful flame retardant and smoke suppressant properties to PVC, polychloroprene, thermoplastic substrates, and chlorowaxes. An addition of 1-5 or up to 10 parts of ZS or ZHS per 100 parts of organic compound significantly reduces the average and the peak rate of heat generation, and also gives a smoke reduction of about 50%. These compounds have also shown promising results as smoke suppressants and flame retardant synergists in halogen-free plastics and elastomers. An addition of ZS or ZHS to alkyd resin-based glass paints improves the flame resistant properties. Tin fluoborate and stannous sulphate are used as electrolytes in the electroplating of tin in the form of tin plate and also as tin alloy coatings.
Tin acetate, as well as tin formate, is used as a catalyst in the hydrogenation of coal and liquid fuels. A considerable improvement in the uptake of dyes in textiles can be achieved by the use of tin acetates or tin tartarate. Tin oxalate is an excellent catalyst for esterification reactions, it limits the extent of detrimental reactions that may otherwise occur during esterification reactions. Stannous octoate is extensively used in the production of polyurethane foams and certain silicone elastomers which can be vulcanised at room temperature. Stannous stearate or stannous oleate are used as additives in polyvinyl chloride (PVC) to improve its stability on exposure to heat and/or light. The presence of the tin compound prevents the removal of HC1 from the PVC. It is the removal of HC1 that is in fact responsible for the poor heat and light stability of PVC. 1.8.3.3 Organotin Compounds The commercial use of organotin compounds is either as tri-organotin or as diorganotin compounds. Tetra-organotin compounds have a limited use in industry but they are quite important because they are used to synthesise practically all the other organotin compounds. One use of tetra-organotin compounds is as a corrosion inhibiting additive in hydrocarbon lubricating oils. Most of the tri-organotins are biocides, these are therefore used for a variety of biocidal activities. Trimethyl tin and triethyl tin are highly toxic to mammals and other life forms, thus their use is very much limited. Tributyl tin and triphenyl tin compounds are very effective against fungi, germs, insects, bacteria, etc., and they have a very low mammalian toxicity. These compounds are therefore extensively used in agriculture for the pro tec-
1.9 References
tion of crops, as additives in cable sheathing plastics for repelling rodents, as additives in anti-fouling paints, in formulations for use in the protection of various marine parts, wood, timber, etc. Some tributyl tin compounds are also used as additives on polymeric molecules such as polysiloxane compounds. These and some other such compounds have been developed as antifoulant polymers for the controlled release of the organotin compound to prevent fouling over extended periods. Di-organotin compounds are chemically more reactive than tri-organotin compounds, but they have a very low influence on biological activity. Most of the di-organotin compounds such as dibutyl tin dilaurate, dibutyl tin, and di-«-octyl tin maleate are used as stabilisers in polyvinyl chloride, as catalysts in the manufacture of rigid or flexible polyurethane foams, and as additives in silicone elastomers for facilitating curing at low temperatures. Dimethyl tin dichloride is a preferred precursor for providing a coating of tin oxide on molten glass.
1.9 References Barry, B. T. K., Thwaites C. I (1983), Tin and its Alloys and Compounds. Chichester, U.K.: Ellis Harwood. Carlin, J. F. (1985), Tin, Mineral Facts and Problems, U.S. Bureau of Mines Bulletin 675. Washington, DC: U.S. Dept of Interior, p. 847. Coles, S. V. G. (1988), "The Use of Tin(iv) Oxide as Selective Gas Sensing Materials", Tin and Its Uses 158, 1. Denholm, W. T., Foo, K. A., Pocock, N. J. B. (1984), in: Proc. CSMjInst. Min. Met. Joint Conf. on Mineral Processing and Extractive Metallurgy, Kunming, China, Oct. 27-Nov. 3, pp. 275-291. Evans, C. X, Karpel S. (1985), Organotin Compounds in Modern Technology. New York: Elsevier. Fletcher, A. W., Jackson, D. V., Valentine A. G. (1967), Trans-Inst. Min. Metall. C 76, C 145-153.
69
Flett, D. S.? Holt, G., Voyzey, R. B. G., Chaston, I. R. M. (1984), in: Proc. CSM/Inst. Min. Met. Joint Conf. on Mineral Processing and Extractive Metallurgy, Kunming, China, Oct. 27-Nov. 3, pp. 581587. Floyd, J. M. (1980), in: Proc. Lead Zinc Tin 80: Cigan, X M., Mackay, T. S., O'Keefe, T. X (Eds.). Feb. 24-28, Las Vegas, NV. Warrendale, PA: TMSAIME, p. 508. Foo, K. A., Lightfoot, B. W. (1987), in: Proc. Pyromet 87, Sept. 21-23. London: Institution of Mining and Metallurgy, p. 389. Germain, W. (1983), in: Encyclopedia of Chemical Technology, Vol. 23: Mark, H. F , Othmer, D. R, Overberger C G., Seaborg, G. T. (Eds.). New York, Wiley, pp. 35-42. Gitlitz, M. H., Moran, M. K. (1983), in: Encyclopedia of Chemical Technology, Vol. 23: Mark, H. R, Othmer D. R, Overberger, C. G., Seaborg, G. T. (Eds.). New York: Wiley, pp. 42-77. Halsall, P. (1984), in: Proc. CSM/Inst. Min. Met. Joint Conf. on Mineral Processing and Extractive Metallurgy, Kunming, China, Oct. 27-Nov. 3, pp. 303-323. Hampshire, W. B. (1992), in: Encyclopedia of Physical Science, Vol. 16. New York: Academic Press, pp. 791-805. Hansen, M., Anderko, K. (1958), Constitution of Binary Alloys. New York: McGraw-Hill, p. 52. Hasouna, A. T., Nogi, K., Ogino, K. (1988), Trans. Jpn. Inst. Met. 29, 748-755. Hedges, E. S., Cuthbertson, X W, Ellwood, E. C , Hoare, WE., Lewis, W R. (1960), Tin and its Alloys. London: Edward Arnold. Karpel, S. (1986), "Gas Sensors Based on Tin Oxide", Tin Its Uses 149, 1. Luhman, X, Dew Hughes, D. (1979), Metallurgy of Superconducting Materials; Treatise on Materials Science and Technology, Vol. 14, New York: Academic Press. Mackay, C. A. (1977), Aspects of Soldering, Int. Tin Res. Inst. Publ. No. 539. Uxbridge, England: International Tin Research Institute. Manko, H. (1964), Solder and Soldering. New York: McGraw-Hill. Massalski, T. B. (1986), Binary Alloy Phase Diagrams. Metals Park, OH: American Society for Metals. Maykuth, D. X (1983), in: Encyclopedia of Chemical Technology, Vol 23: Mark, H. R, Othmer, D. R, Overberger, C. G., Seaborg, G. T. (Eds.). New York: Wiley, pp. 18-35. Miller, D. R. (1989). "Tin, a Vital Component of Dental Amalgam", Tin Its Uses 160, 1. Pearson, D., Holt, G., Winter, D. G. (1977a), Trans.Inst. Min. Metall. C 86, 140-146. Pearson, D., Holt, G., Winter, D. G. (1977b), Trans.Inst. Min. Metall. C 86, 175-185. Pitman, K. (1988), Min. J (London) 37.
70
1 Tin
Villars, P., Calvert, L. D. (1985), Pearson's Handbook of Crystallographic Data for Intermetallic Phases. Metals Park, OH: American Society for Metals. Wadekar, S. L., Banerjee, S., Raman, V. V., Asundi, M. K: (1991), Zirconium in Nuclear Industry: 9th Int. Symp., ASTM STP 1132: Eucken, C. M., Garde, A. M. (Eds.). Philadelphia, PA: ASTM, pp. 140-155. Wright, P. A. (1983), Extractive Metallurgy of Tin, 2nd ed. New York: Elsevier. Yokota, M., Nose, M., Takano, Y, Mitani, H. (1980), J. Jpn. Inst. Met. 44, 170.
General Reading Barry, B. T. K., Thwaites, C. J. (1983), Tin and Its Alloys and Compounds. Chichester, U.K.: Ellis Harwood. Belyayov, N. M. (1963), A Handbook of the Metallurgy of Tin. Oxford: Pergamon. Hampshire, W. B. (1992), Encyclopedia of Physical Science, Vol. 16. New York: Academic. Hedges, E. S., Cuthbertson, J. W., Ellwood, E. C , Hoare, W. E., Lewis, W. R. (1960), Tin and Its Alloys. London: Edward Arnold. Mantell, C. L. (1970), Tin - Its Making, Production Technology and Applications. New York: Hafner. Mark, H. R, Othmer, D. E, Overberger, C. G., Seaborg, G. T. (Eds.), (1983), Encyclopedia of Chemical Technology. New York: Wiley. Wright, P. A. (1983), Extractive Metallurgy of Tin, 2nd ed. New York: Elsevier.
2 Lead Alloys R. David Prengaman RSR Corporation, Dallas, TX, U.S.A.
List of 2.1 2.1.1 2.1.2 2.1.3 2.1.4 2.1.5 2.1.6 2.1.7 2.2 2.2.1 2.2.2 2.2.3 2.2.4 2.2.5 2.2.6 2.2.7 2.3 2.3.1 2.3.2 2.3.3 2.3.4 2.3.5 2.3.5.1 2.3.5.2 2.3.5.3 2.3.5.4 2.3.5.5 2.4 2.4.1 2.4.2 2.5 2.6 2.7
Symbols and Abbreviations Introduction Major Uses Environmental Aspects of Lead Usage Recycling Physical Properties Unique Characteristics of Lead Electrochemical and Chemical Properties Alloying Behavior Lead-Antimony Alloys Mechanical Properties of Cast Lead-Antimony Alloys Cast Lead-Antimony Alloys in Lead-Acid Batteries Cast Low-Antimony Alloys for Batteries Ammunition Anode Alloys Other Uses of Cast Lead-Antimony Alloys Wrought Lead-Antimony Alloys Lead-Calcium Alloys Uses of Lead-Calcium Alloys Cast Lead-Calcium Alloys Strengthening Mechanism in Cast Lead-Calcium Alloys Wrought Lead-Calcium Alloys Lead-Calcium-Tin Alloys Cast Lead-Calcium-Tin Alloys Electrochemical Effects of Tin Wrought Lead-Calcium-Tin Alloys Additives to Lead-Calcium-Tin Alloys Reactivity of Lead-Calcium-Tin Alloys Lead-Tin Alloys Solders Corrosion Protection Other Lead Alloys Concluding Remarks References
Materials Science and Technology Copyright © WILEY-VCH Verlag GmbH & Co KGaA. All rights reserved.
72 73 73 73 74 74 75 75 76 76 77 78 79 80 80 80 81 81 82 82 82 83 83 83 84 84 85 85 85 86 86 87 87 87
72
2 Lead Alloys
List of Symbols and Abbreviations a r
lattice parameter ratio of tin to calcium in P b - C a - S n alloys
f.c.c. BCI IACS IBMA ILZRO ILZSG OSHA SLI SR UTS YS
face centered cubic Battery Council International international annealed copper standard Independent Battery Manufacturers Association International Lead-Zinc Research Organization International Lead/Zinc Study Group Occupational Safety and Health Agency (U.S. Government) starting and ignition battery stress rupture, time (h) to failure at designated stress ultimate tensile strength (MPa) yield strength (MPa)
Compositions are given in weight percentages unless otherwise stated.
73
2.1 Introduction
2.1 Introduction In 1993, the estimated world consumption of lead was 4.47 Mt. The world consumption has stagnated 1989. (See also the Introduction to the volume.) Deceased consumption in Eastern Europe and the CIS countries has been offset by increased consumption in Southeast Asia and the U.S. (ILZSG, 1994). Some of the major uses for lead are decreasing rapidly owing to environmental considerations. This decrease has offset increased lead usage for lead-acid storage batteries. Batteries have experienced annual growth in excess of 3%. Lead consumption is expected to grow at a rate of 2 - 3 % per year as the Eastern Europe economies are stabilized and the dissipative lead usages such as antiknock compounds are eliminated. About half of the lead is produced from ore (galena) and half from recycled scrap, primarily used automobile batteries and building materials. About half of the lead is used in the form of alloys and the remainder as lead oxides for the active materials of lead-acid batteries and ceramics. 2.1.1 Major Uses The major use for lead worldwide is lead-acid batteries. Table 2-1 shows the lead usage in the USA in 1963 and 1992 (U.S. Bureau of Mines, 1963, 1993). Many of the applications listed in 1963 no longer remain. Other uses are ammunition, brass and bronze, cable coverings, solders, sheet lead, anodes, pigments, oxides for glass manufacture, weights, and ballast. In the developed countries, lead usage other than lead-acid batteries and several other applications is declining due to the toxic nature of lead.
Table 2-1. Consumption of lead (tons) in U.S. by product. Product
1992
1963
Ammunition Bearing metals Brass & bronze Cable & caulking Pipe & tube Sheet Solder Casting metals Storage batteries Oxides & pigments Gasoline & additives Miscellaneous
64 800 4 800 9 200 17 000 11700 21000 13 500 20100 998 200 63 200 3 000 10 000
45 400 19 700 20 000 121 800 35 400 24100 61800 32 600 399 200 90100 175 900 31800
1 236 500
1 057 700
2.1.2 Environmental Aspects of Lead Usage Lead is detectable in almost all phases of the inert environment and in all biological systems. Lead is toxic to most living things at high exposure. In the toxicology of lead in humans, the major risk is toxicity to the central nervous system (Rutter and Jones, 1983; Niragu, J., 1978). Nearly all environmental exposure to lead is via inorganic compounds. The principal route of exposure is food or water, however, other exposures such as leadbased paint in older dwellings, lead in air from combustion of lead containing auto exhausts, industrial emissions and handto-mouth activities of young children may also contribute to lead exposure. Blood lead levels are a good indicator of recent exposure to lead. The body burden is stored in the skeleton and soft tissues. The greatest interest at the present time concerns the maximum level of lead exposure in the neonatal and young child which does not produce a cognitive or neurological effect (Gower, 1990). A second concern
74
2 Lead Alloys
is the industrial exposure of adults. OSHA has set acceptable levels of exposure and blood lead levels for adults working in industries where they may be exposed to lead (OSHA, 1978). Because of the environmental concerns about lead, many lead uses have been banned and the application of lead will continue to be restricted. Many lead products such as lead-based paints, lead-antiknock compounds, lead solder for water pipes, lead glazes for pottery and dinnerware, lead shot for waterfowl hunting, lead collapsible tubes, lead foil for sealing wine bottles, and water and sanitary pipes have been eliminated. In the western world, the dissipative uses for lead are rapidly being discontinued, while those uses where the lead alloys can be recycled remain and may continue to grow. 2.1.3 Recycling Lead is the most highly recycled metal. In 1990, recycled lead accounted for 2.3 million tons while primary production accounted for 2.06 million tons. Thus, recycled lead constituted 52.6% of total lead production worldwide compared to 34.8% in 1970 (Boehnke, 1991). In other countries such as the United States, recycled lead in 1992 constituted 75% of lead poduction and consumption (U.S. Bureau of Mines, 1993). The increase in lead recycling is the result of concentration of lead in the recyclable products such as batteries, shielding, anodes, flashings and weatherings and cable coverings. The dissipative uses for lead such as paint, antiknock compounds, ammunition, solder water and waste pipes have been banned or are being slowly eliminated. In the U.S., batteries were recycled at a rate of 97% in 1992 (Battery Council International, 1993).
Since batteries constitute the major raw material for recycling, processes have been developed (Olper and Asano, 1989) to break batteries for recycling. The major lead recycling companies in the world are RSR Corporation (U.S.), Exide Corporation (U.S.), Sanders Lead Company (U.S.), Metallgesellschaft (Europe), Metaleurop (Europe), Samim (Italy), Schuylkill Metals (U.S.), H. J. Enthoven Limited (U.K.). The major battery recycling processes are blast furnaces (Pike, 1990), rotary furnaces (Forrest and Wilson, 1980), reverberatory furnaces (Prengaman, 1991a) and electric furnaces (Eby, 1990). These are pyrometallurgical processes. New hydrometallurgical processes to electrowin lead from battery pastes have been developed but have not yet been implemented (Prengaman and McDonald, 1990; Olper, 1991; Serracane, 1991; Cole and Lee, 1983). In the future, lead products will have the requirements for mandatory recycling which should increase the recycling rate worldwide to 65-75%. 2.1.4 Physical Properties The physical properties of pure lead are shown in Table 2-2 (Worchester and O'Reilly, 1990). Lead is a bluish grey metal. It has a low melting point (327.3 °C), high boiling point, and low heat of fusion. This makes it an ideal material as a casting metal. Lead is relatively fluid and because of the low melting point can be cast into a variety of shapes using many mold materials. The high fluidity permits the use of lead and lead alloys for intricate metal shapes as well as thin battery grids (< 1 mm thick). The low melting point and high density presents a problem because they limit the mechanical properties even at room temperature. Because of its low modulus of elasticity and f.c.c. crystal structure, lead
2.1 Introduction
Table 2-2. Physical properties of lead. Property Crystal structure Lattice parameter Atomic weight Density at 20 °C Melting point Boiling point Volume contraction upon solidification Mean specific heat (100-200 °C) Thermal conductivity 25 °C Thermal expansion 20-100°C Latent heat of fusion Electrical conductivity Electrical resistivity (20°C) Modulus of elasticity (20 °C)
Unit
A g/cm3 °C °C % J/(kgK)
Value f.c.c a = 4.94 207.21 11.34 327.3 1740 3.5 129
35 W/(m K) um/(m K) 29.3 xl0~ 6 kJ/kg %/IACS nQm
GPa
23 8.3 206.43 14
and lead alloys are readily formed into useful shapes by rolling, extruding, forging, and spinning at room temperature. The low modulus of elasticity, low recrystallization temperature, and relatively high coefficient of expansion, combined with the high density limit the use of lead as a structural material without alloying elements or proper support. The high density and low elastic limit makes lead an excellent barrier material for reducing the transmission of noise from adjacent areas. The high density of lead (11.35 g/cm3) is also ideal for the attenuation of radiation. Lead is used extensively for shielding of x-rays and gamma-rays. The excellent malleability and low melting point make possible the production of complex shielding. 2.1.5 Unique Characteristics of Lead
Lead is a unique material in that it is used both as a structural and corrosion resistant material as well as the active ma-
75
terial in lead-acid batteries which undergoes oxidation and reduction during the charge-discharge reactions. The ability of lead to exist in the 0, + 2, and + 4 valence states makes lead-acid batteries possible. Lead alloys serve as the structure which holds the lead oxides of the active material. The grid structure must be strong (YS 40 MPa) (Perkins, et al., 1976), resist corrosion during the oxidation conditions of charging, and be conductive to carry high currents from the battery. The purity of lead used for the active material of lead-acid batteries is very important. The impurity limits of pure lead in many countries of the world is shown in Table 2-3. The major detrimental impurities affecting battery behavior are Sb, As, Cu, Ni, Te, Fe, and Zn. Purity of lead can reach 99.99% or higher. Bismuth is the major impurity which dictates overall lead content. Only the U.S. points to Te and Ni in the specification. These are major gas producing elements and are detrimental to sealed lead-acid batteries. Customer's specifications throughout the world are more stringent than national standards. The permissible levels for impurities in sealed maintenance-free lead-acid batteries has been defined (Pierson et al., 1975). 2.1.6 Electrochemical and Chemical Properties
Lead and many of its alloys have excellent corrosion resistance to a variety of corrosive substances due to the formation of a passive, impermeable, insoluble protective film when the lead is exposed to the corrosive solution. The passive film protects the surface of the lead unless the film is damaged or destroyed. Lead shows excellent resistance to H 2 SO 4 and H 3 PO 4 at all concentrations
76
2 Lead Alloys
Table 2-3. National standards for trace impurities in pure lead for batteries (wt-%). Element
Ag As Bi Cu Fe Ni Sb Sn Te Zn Cd
DIN1719 Germany
UN 13165 Italy
ASTM B29-93 US
99.95 Australia
HP2 Canada
BS334 UK
0.001 0.001 0.010 0.001 0.001
0.0015 0.002 0.010 0.0015 0.0025
0.002 trace 0.005 0.003 0.003
0.0035 0.001
0.002 0.002 0.050 0.002 0.002 0.002 0.002 0.002
0.0015 0.0015 0.005 0.0015 0.002
0.001 0.001
0.0015 0.0015
0.002 trace
0.001
0.002 0.001
0.0025 0.0005 0.025 0.001 0.001 0.0002 0.0005 0.0005 0.0001 0.0005 0.0005
0.002 0.002
0.001
0.002 trace
as well as sulphides, sulphites, and sulphate salt solutions. The corrosion film is insoluble lead sulphate. Lead is not resistant to nitric or acetic acid because the reaction products are soluble permitting continuous attack. In most solutions, lead and lead alloys are resistant to galvanic corrosion because of the formation of non-conductive films. In contact with noble metals, however, lead can undergo galvanic attack (Smith, 1979). Lead alloys are used as insoluble anodes in electrowinning of primary metals (Zn, Ni, and Cu) as well as anodes for MnO 2 production, chrome plating, and galvanic protection using applied currents (Prengaman, 1984 b). In the plating solutions, the anodes form an insoluble PbO 2 layer during use. The PbO 2 layer is hard and adheres to the lead alloy substrate to prevent corrosion and prevent contamination of the cathode with lead. Alloys of lead combined with Ag, Sn, Sb, Ca, and As are used for electrowinning anodes. 2.1.7 Alloying Behavior
Lead forms alloys with a wide variety of metals. The principal commercial lead al-
loys contain Sb, Sn, and Ca. Minor alloying additions are Cu, As, Ag, Se, and S. The solubility of most of the elements is high but declines at room temperature leading to strengthening by precipitation. A listing of the various alloys of lead can be found in the Metals Handbook (Worchester and O'Reilly, 1990).
2.2 Lead-Antimony Alloys Pb-Sb alloys are the most widely used lead alloys. Pb-Sb is the primary alloy used in lead-acid batteries througout the world. Alloys are also used for ammunition, cable sheathing, anodes, bearing metals, solders, special casting metals, and wrought products. The phase diagram for the Pb-Sb system is shown in Fig. 2-1 (Hansen, 1958). The Pb-Sb system is an eutectic system with the eutectic point at 11.1% Sb and 252 °C. (All percentages are by weight.) The solubility of Sb in lead decreases from 3.5 wt.% at 252°C to 0.25 wt.% at 25 °C Thus, the alloys are hardenable via precipitation of the Sb. Virtually, all the commercial Pb-Sb alloys used today are the hy-
2.2 Lead-Antimony Alloys
11
Antimony, at. %
700
10
20
40
50
60
80
70
90
l
600 -
Pb
30
Liquid
10
30
40
70
50 60 Antimony, wt %
80
90
Sb
Figure 2-1. Lead-antimony phase diagram (Hansen, 1958).
poeutectic alloys. Alloys with Sb contents lower than 3.45% Sb should contain no eutectic; however, as a result of coring, eutectic is normally present to 1% Sb and segregation of Sb can be seen in grain structures below that value. Most Pb-Sb alloys are used in the cast form. The mechanical properties of cast P b Sb alloys, particularly, those in the 0.55.0% range, depend to a high degree on the thermal history of the material. The nature of the pre-treatment can result in differences of over 100% in the resultant mechanical properties. Other alloying elements such as As, Cu, Sn, S, and Se can also have an effect on the mechanical properties as can the solidification conditions. The age-hardening response corresponds to the precipitation of particles of Sb on the {111} planes of the lead f.c.c. crystals (Aballe et al, 1972).
2.2.1 Mechanical Properties of Cast Lead-Antimony Alloys
The mechanical properties of Pb-Sb alloys are influenced by the Sb content, rate of solidification, thermal treatments, and aging time. Table 2-4 shows the effect of Sb content and aging time on the mechanical properties of cast Pb-Sb alloys (Prengaman, 1983). High-Sb alloys (> 6% Sb) are strengthened primarily by the Sb particles of the eutectic phase produced during solidification of the alloy. Lower-Sb alloys are strengthened by a combination of eutectic and precipitated Sb particles. In Pb-Sb alloys, precipitation-hardening occurs rapidly, particularly when the alloys are quenched immediately after casting from elevated temperatures. Table 2-5 (Dean et al., 1926) shows the effect of
78
2 Lead Alloys
Table 2-4. Mechanical properties of aged lead-anti-
mony alloys. Antimony content (wt.%)
11 6 3 2 1 0
YS (MPa) Days after aging at 25' C
UTS (MPa)
Elongation (%)
1 day
30 days
30 days
30 days
68.9 55.2 34.0 24.1 13.8 3.5
74.4 71.0 55.2 37.9 19.3 3.5
75.9 73.8 65.5 46.9 37.9 11.7
5 8 10 15 20 55
Table 2-5. Effect of quenching temperature on the strength of lead-3% antimony, aged 1 day at 25 °C. Quenching temperature
Tensile strength (MPa)
238 230 215 198
75 56 38 34
quenching temperature on the properties oflead-3% Sb alloys. Minor alloy additions of As and Cu result in significant increases in the mechanical properties of Pb-Sb alloys and may change the precipitation kinetics by changing the morphology of the precipitate (Bluth and Hanemann, 1937; Mao and Larson, 1968; Mao et al., 1969; Heubner and Ueberschaer, 1974). The addition of As also can change the mode of the precipitation reaction from continuous precipitation to a discontinuous reaction at low (<2%) Sb contents (Borchers and Reuleaux, 1980).
2.2.2 Cast Lead-Antimony Alloys in Lead-Acid Batteries
Pb-Sb alloys have been used as electrodes, connectors, and terminals of leadacid batteries since the 1880s (Wade, 1902; Vinal, 1955). Pb-Sb alloys are strong, creep-resistant and corrosion-resistant in H 2 SO 4 , and thus resist structural changes caused by charge-discharge cycles in the battery. Sb as an alloying element in the positive electrode grid aids in retaining the active material in the battery and also aids in recharge of the battery when it is deeply discharged. Deep-discharge industrial, load-leveling, and motive-power batteries utilize lead alloys containing 5 - 1 1 % Sb. Large tubular grids use 9 - 1 1 % Sb. This alloy freezes at the eutectic temperature and is capable of flowing long distances before freezing. Optimum castability of Pb-Sb alloys has been demonstrated to be near the eutectic composition (White and Rogers, 1974). Lead-acid batteries produced from P b Sb positive grids have one disadvantage. During use, the positive grid is corroded during charging. The Sb oxidized from the positive grid enters the electrolyte and plates on the cathode. Sb at the cathode reduces the voltage at which water in the electrolyte breaks down into H + and O 2 ~. During charging of the battery water is consumed by the gassing and must be replaced (Ruetschi and Cahan, 1975; Dawson et al., 1970). The amount of Sb transferred increases with service life and Sb contant of the positive grid. There are significant electrochemical benefits to the presence of Sb in the grid alloy related to positive electrode performance (Burbank, 1964). Sb seems to improve the adhesion of the active material to the grid as well as improve its electrochemical activity (Simon, 1965).
2.2 Lead-Antimony Alloys
79
2.2.3 Cast Low-Antimony Alloys for Batteries Reducing the Sb content of the positive grid reduces the migration of Sb. Batteries using Pb-Sb alloys containing 1.5-2.75% Sb experience reduced gassing. In this alloy range, however, the (liquid+ Pb) range of the phase diagram and the freezing range are increased. Only a small amount of eutectic liquid remain at the final freezing temperature. Low-Sb alloys solidify into a coarse dendritic structure which gives rise to hot tearing, coarse grains, and brittleness in casting. In battery grids, porosity and cracks are accentuated near grid wire intersections, where shrinkage stresses are highest (Heubner and Ueberschaer, 1974). Figure 2-2 shows the structure of a Pb-Sb alloy with a coarse dendritic structure. To prevent the cracking, small amounts of S, Cu, or Se are added to serve as sites for growth of lead crystals during solidification to produce fine-grained, pore-free castings resistant to cracking (Heubner etal, 1975; Prengaman 1977a, 1983; Kallup and Berndt, 1984; Kallup, 1986). Figure 2-3 shows the grain structure of a Pb-Sb alloy modified by nucleants. Cu reacts with As, S, or Sb to produce Cu3As, Cu 2 S, or Cu2Sb. S and Se react with lead to produce PbS or PbSe. Cu and Se are generally used at higher Sb contents while Se is used for lower Pb-Sb alloys. The grain refiners along with As also tend to modify the eutectic phase changing it from an acicular to a more globular structure (Heubner and Sandig, 1971). Sn additions to low-Sb alloys improve fluidity and mold detail, but also tend to negate the effects of the nucleants, producing coarse dendritic structures (Mao et al, 1969). Other alloys have been suggested for low-Sb uses. The major one is a P b - S b Cd alloy (Bagshaw, 1968). This alloy has
Figure 2-2. Grain structure of Pb-3.5% Sb alloy without nucleants. Magnification x 320.
Figure 2-3. Grain structure of Pb-3.5% Sb alloy with nucleants. Magnification x 320.
80
2 Lead Alloys
many of the castability aspects of higher Sb alloys but with a dramatic reduction of Sb transfer to the negative grid. Ag has been shown to dramatically increase the creep resistance and UTS of Pb-Sb alloys and to increase the conductivity of the corrosion layer. Pb 3-4.5% Sb alloys are also used for the terminals and internal cell connectors in lead-acid batteries. The cast alloy has high mechanical properties and wets the grid lugs with eutectic liquid to form a complete metallurgical bond (Prengaman, 1989).
regular surfaces. The anodes develop a hard PbO 2 corrosion layer which makes them resistant to corrosion. Pb-Sb anodes resist passivation when the current is interrupted and the anode may remain in the plating solution for long periods of time without damage. Pb 6-10% Sb anodes containing 0.51.0% Ag are used as anodes for cathodic protection of pipelines, oil-drillling platforms, and ocean-going vessels under an applied current. The Ag addition reduces corrosion from chloride-containing environments.
2.2.4 Ammunition
2.2.6 Other Uses of Cast Lead-Antimony Alloys
Pb-Sb alloys are also used in the cast form for ammunition. Lead shot is produced by dropping molten Pb-Sb alloys containing 0.5-8% Sb through small holes drilled into iron pans. The molten drops freefall through the air and are quenched in water at the bottom of the tower. The alloys contain As in an amount equal to about 25-30% of the Sb content. As dramatically increases the surface tension of the P b - S b - A s alloy, permitting the molten droplet to attain a round shape during freefall. Sn and Cd, if present, reduce the surface tension and produce elongated shot, and are generally restricted to less than 5 ppm. Shot up to 5.8 mm is dropped from towers, larger shot is cast into molds. 2.2.5 Anode Alloys
Pb - Sb alloys containing 0.5-1 % As have been used in the cast form as anodes for Cu, Ni, and Cr electrowinning and plating processes (Taiton, 1928). The Pb-Sb anodes have sufficient mechanical properties to span plating tanks and can be cast into intricate shapes to accurately plate ir-
P b - S b - S n alloys contain high amounts of antimony (10-16% Sb) and 1-7% Sn are used for printing alloys, bearings, solders, and specialty castings. The alloys have very low melting points, high hardness, excellent fluidity, and produce excellent surface finish. Pb-Sb printing alloys were used for the first moveable type but have been replaced by electronic typesetting. Excellent tribological properties and good hardness makes P b - S b - S n alloys suitable for journal bearings. The alloys contain 9-15% Sb and 1-20% Sn and may also contain Cu and As, which improve compression, fatigue, and creep strength important in bearings. Pb-Sb-Sn bearing alloys are listed in ASTM B23-83 (ASTM, 1993). Specialty castings include belt buckles, trophies, casket (coffin) trim, miniature figures, and hollow ware. Slush castings are produced by pouring an alloy into a mold, permitting a given thickness to solidify, and pouring out the remaining liquid to produce a hollow, light, high-detail casting. The near-eutectic lead alloy, contain-
2.3 Lead-Calcium Alloys
ing 11% Sb, 1% Sn, and 0-0.5 As, has a low melting point which permits silicon rubber molds to be used in spin-casting processes. Low Sb (2-5%), low Sn (2-5%) lead alloys are used for automobile body solder. Special Pb-Sb alloys containing 1-4% Sb are used for wheel-balancing weights, battery cable clamps, and highly machined isotope pots. 2.2.7 Wrought Lead-Antimony Alloys
Wrought Pb-Sb in the form of rolled or extruded products is significantly weaker and has much higher creep rates than the cast material (Borchers and Scharfenberger, 1966). Table 2-6 shows the mechanical properties of cast and rolled P b - 6 % Sb alloy. During deformation, the eutectic strengthening network is destroyed and broken up into smaller pieces. The alloy is subjected to recrystallization since most rolling is performed at elevated temperatures to prevent cracking. Pb-Sb rolled and extruded products have been replaced in the chemical industry by stainless steel and plastic, but still find extensive use in H 2 SO 4 production systems as well as acid and sulfate heating and cooling coils. As and Cu are added to the alloys to increase creep resistance at elevated temperatures. Because of the low mechanical properties, rolled Pb-Sb alloys have not been used as alloys in continuous lead-acid battery grid production. Pb 0.5-1.0% Sb alloys have been used for the protective sheaths over telephone and fiber optic cables. These alloys are generally stronger and more creep resistance than Pb-Cu, Pb-Sn, and P b - S n - A s alloys also used for cable coverings. Pb-Sb alloys with less than 0.6% Sb can be continuously extruded, are impervious to
81
moisture and oil, and remain pliable indefinitely. They do not suffer the age-hardening of other alloys of the same strength.
2.3 Lead-Calcium Alloys Ca alloys readily into lead. The lead-rich portion of the P b - C a phase diagram is shown in Figure 2-4 (Hansen, 1958). The P b - C a system has a peritectic with the peritectic temperature of 328.3 °C, an elevation of 1°C from the melting point of pure lead of 327.3 °C. Cast alloys above 0.07% Ca show primary crystallites of Pb 3 Ca. Samples containing less than 0.10% calcium held at a
Table 2-6. Mechanical properties of cast and rolled P b - 6 % Sb alloy.
UTS (MPa) YS(MPa) Elong. (%) Stress rupture (hours at 20.7 MPa)
Cast
Rolled
73.8 71.0 8 1,000
30.6 19.5 35 1.5
Calcium, at.% 0.2 1 Liquid
400
0.4 !
0.6
1 •
300
_327.3 # C a
^0.1056
0.8
1.0
i 1 a + liquid
3283 # C _
jr
-
200
t+0 100 -
/
L—0.01% 0 Pb
0.04
1 0.08 0.12 Calcium, wt %
I 0.16
0.20
Figure 2-4. Lead-calcium phase diagram (Hansen, 1958).
82
2 Lead Alloys
temperature of 326 °C show disintegration of the primary crystallites as would be expected in a peritectic reaction. The solubility of calcium in lead decreases from 0.10% at the peritectic temperature to about 0.01% at25°C. Only the phase Pb 3 Ca is of interest in P b - C a alloys. Pb 3 Ca has an f.c.c. structure with calcium ions at the cube corners and lead ions at the face centers (Zintl and Neumayr, 1933). The density of Pb 3 Ca is 9.4 g/cm 3 . The lattice parameter a = 4.901 A is about 1 % smaller than that of lead. Solid solutions of calcium in lead precipitate Pb 3 Ca during subsequent aging, resulting in significant increases in strength. 2.3.1 Uses of Lead-Calcium Alloys P b - C a alloys were initially developed as alloys for cable sheathing (Schumacher and Buton 1930; Dean and Ryjord, 1930); sheet (Greenwood and Cole, 1949) and storage batteries (Schumacher and Phipps, 1935; Haring and Thomas, 1935). P b - C a alloys are used extensively in standby power batteries, for telephone service, submarine batteries, and the negative grid of SLI hybrid batteries (Prengaman, 1992 a). P b - C a alloys have also been used as alloys for bearings. The higher float voltage (the voltage at which standby power batteries are continuously charged to permit instant availability to the circuit) and higher hydrogen overvoltage for P b - C a alloy batteries has reduced maintenance and eliminated the possible H3As or H 3 Sb generation in submarine and standby power batteries (Willingantz, 1972). 2.3.2 Cast Lead-Calcium Alloys
The major use of P b - C a alloys is the production of grids for lead-acid storage batteries. The excellent fluidity of P b - C a
Table 2-7. Mechanical properties of cast Pb-Ca alloys. Calcium content (wt. %) 0.03 0.04 0.06 0.08 0.10 0.12 0.14
UTS (MPa)
Elongation (%)
Stress rupture (hours at 20.7 MPA)
30.7 36.8 42.9 46.4 47.8 43.2 39.2
40 40 40 30 35 40 40
1 10 40 40 10 5 2
alloys permits extremely thin grids to be cast successfully. The high freezing point (328.3 °C) permits rapid solidification and ejection from the mold at high temperatures. The final mechanical properties depend on the cooling rate from room temperature and the calcium content. The mechanical properties of fully-aged binary P b - C a alloys are shown in Table 2-7 (compiled from data developed by Rose and Young, 1974; Townsend, 1960; Prengaman, 1987). 2.3.3 Strengthening Mechanism in Cast Lead-Calcium Alloys
P b - C a alloys are strengthened by precipitation of Ca in the form of Pb 3 Ca from solid solution. As the temperature of the castings is reduced, the solubility of the Ca is reduced. The age-hardening of P b - C a alloys is very pronounced at Ca contents as low as 0.03% calcium. Because Ca has a much larger atomic radius than lead, diffusion is difficult, and precipitation in P b - C a alloys occurs via a discontinuous precipitation mechanism (Scharfenberger and Henkel, 1973; Borchers et al, 1975; Prengaman, 1976; Caillierie et al., 1986). The discontinuous precipitation reaction is similar to that in
2.3 Lead-Calcium Alloys
Sn (Tu, 1972; Tu et al., 1969, 1967). The detailed hardening mechanism is complex and involves three successive discontinuous reactions to the hardening process (Bouirden et al., 1991) which can take as long as 30 days. In the discontinuous precipitation process grain boundaries move into the supersaturated solid solution and the precipitate Pb 3 Ca particles in a lamellar array as they move through the material. During solidification of P b - C a alloys, Ca is segregated to the interdendritic regions of the casting. Bi increases the segregation and the rate of discontinuous precipitation reaction (Prengaman, 1993; Caldwell and Sokolov, 1976). The discontinuous precipitation reaction may begin at the interdendritic boundaries, in many locations at the same time in the casting, and gives the classic serrated grain structure of P b - C a alloys which is shown in Figure 2-5. 2.3.4 Wrought Lead-Calcium Alloys
The early investigations of P b - C a alloys involved working the castings to produce extruded or rolled forms. Extruded Pb-Ca alloys are significantly more creep
Figure 2-5. Grain structure of Pb-0.082% Ca alloy: serrated grain boundaries. Magnification x 160.
83
and fatigue resistant than other alloys used for cable shearing (Dean and Ryjord, 1930; Greenwood and Cole, 1949). The energy of the rolling process is dissipated by the discontinuous precipitation process and overaging of the wrought products occur (Borchers et al., 1975; Prengaman, 1980). Deformation of high calcium alloys to produce high-strength P b - C a alloys has been reported (Baker, 1993). 2.3.5 Lead-Calcium-Tin Alloys
P b - C a - S n alloys are more important than P b - C a alloys. Sn increases the mechanical properties of the alloy, particularly when combined with deformation. P b - C a - S n alloys are more corrosion-resistant that P b - C a alloys and at higher Sn contents are not passivated when used as anodes. P b - C a - S n alloys are used as the positive grids for lead-acid batteries, electrowinning anodes, and structural alloys for shielding. 2.3.5.1 Cast Lead-Calcium-Tin Alloys
P b - C a - S n alloys have dramatically differing precipitation characteristics, depending on the Ca and Sn contents (Rose and Young, 1974; Borchers and Assmann, 1978,1979; Prengaman, 1980; Adeva et al., 1980). The structures and aging behavior of continuously cast P b - C a - S n differ from conventional castings (Piercy, 1976). The aging reactions in P b - C a - S n alloys as well as the hardening processes are different for different Sn/Ca ratios. Models to describe the mechanisms for precipitation in P b - C a - S n alloys have been developed (Bouirden et al., 1991; Hertz et al., 1993). Alloys are influenced by the Sn/Ca ratio (r). For low values of r, the hardening process is identical to that in P b - C a alloys. At
84
2 Lead Alloys
high values of r, continuous precipitation of (PbSn)3 Ca occurs followed by several discontinuous precipitation stages. The final intermetallic compound approaches Sn3Ca. The enthalpy of formation of Pb 3 Ca is -34.0 + 0.5 kJ (molatom)" 1 while that for Sn3Ca is -38.6 + 0.8 kJ (mol atom)" 1 at 300°K. Sn3Ca is more stable than Pb 3 Ca and has a higher misfit with the lead lattice. The cellular structure caused by Ca and Sn segregation during solidification accelerates the aging reaction kinetics of the discontinuous reaction. The mechanical properties of fully aged cast P b - C a - S n alloys is shown in Table 2-8 (Rose and Young, 1974). 2.3.5.2 Electrochemical Effects of Tin
Sn additions to P b - C a alloys not only improve the mechanical properties, but also affect the electrochemical behavior (Giess, 1984; Pavlov, 1991; Kostic, 1982). Sn reduces the passivation of lead anodes or positive grids in lead-acid batteries. The reduced tendency toward passivation alows P b - C a batteries with battery grids containing Sn to be deeply cycled. Recent work has shown that the PbO 2 corrosion layer on P b - C a - S n alloys with Sn contents in excess of 1.5% has electronic conductivity in addition to ionic conductivity. Sn also decreases the rate of corrosion P b - C a - S n alloys (Prengaman, 1977b). The decrease in corrosion of cast P b 0.06% Ca-Sn alloys as a function of Sn content is presented in Table 2-9. 2.3.5.3 Wrought Lead-Calcium-Tin Alloys
Deformation combined with the precipitation reaction dramatically effect the properties and grain structures of wrought P b - C a - S n alloys. Rolled P b - C a - S n alloys are used for continuously processed
Table 2-8. Mechanical properties of cast P b - C a - S n alloys. Calcium content (wt. %) 0.06 0.08 0.08 Tin content (wt. %) 0.50 0.50 1.0 UTS(MPa) 46.4 50.0 57.1 Elong. (%) 20.0 35.0 20.0 Stress rupture 150.0 100.0 450.0 (hours at 20.7 MPa)
0.12 0.12 0.5 1.0 50.8 57.1 35.0 30.0 12.0 120.0
Table 2-9. Corrosion of cast 0.065% calcium Pb-CaSn alloys in 1250 SG H 2 SO 4 . Tin content (wt. %)
Corrosion rate (mm/year)
0 0.25 0.50 0.75 1.00 1.25 1.50
0.365 0.345 0.307 0.295 0.289 0.279 0.268
lead-acid battery grids (expanded and punched), as well as high-strength shielding and anodes for copper electrowinning. By combining the effect of deformation and precipitation, extremely fine-grained, corrosion-resistant, high-strength P b - C a Sn wrought materials may be produced. (Prengaman, 1980, 1984 b; Rose and Young, 1974; Barclay et al., 1980). The aging must occur after deformation. If the aging reactions occur prior to deformation, dramatically lower mechanical properties may occur as a result of the discontinuous precipitation processes. The mechanical properties of selected wrought P b - C a - S n alloys are shown in Table 2-10 (Rose and Young, 1974). The mechanical properties of wrought lead-calcium-tin are very stable and resist the normal recrystallization processes at 25 °C because the dislocations serve as sites for the precipitation process and become pinned. Deformation modifies the normal precipitation process in cast alloys
2.4 Lead-Tin Alloys Table 2-10, Mechanical properties of cast P b - C a - S n alloys. Calcium content (wt. %) Tin content (wt. %) Aging before rolling (days) UTS(MPa) Elongation (%) Stress rupture (hours at 27.7 MPa)
0.04 0.50 0 53.5 20.0 12.0
0.06 0.70 0
0.06 1.0 0
0.06 0.06 1.5 1.0 0 0
64.3 72.8 78.5 35.7 15.0 14.0 12.0 35.0 75.0 400.0 1000.0 12.0
by restricting the regions in which the discontinuous precipitation process takes place. 2.3.5.4 Additives to Lead-Calcium-Tin Alloys
Small amounts of other alloying elements such as Bi, Al, and Ag are added to P b - C a - S b alloys. Bi increases the rate of hardening (Myers et al.5 1974) and may increase the resistance to passivation in a manner similar to that of Sn (Pavlov, 1991). Ag has been shown to lead to a marked increase in the creep resistance of cast P b - C a - S n alloys (Bagshaw, 1991, 1989). In addition, Ag increases the conductivity of the corrosion layer, reducing passivation. Al additions to P b - C a - S n alloys reduce the loss of Ca and reduce the amount of suspended CaO in the melt. Al also acts as a grain refiner to reduce the initial as cast grain size of P b - C a - S n alloys and serves to accelerate the first discontinuous precipitation reaction and to delay the second discontinuous precipitation reaction (Prengaman, 1991b; Bouirdenetal., 1991). 2.3.5.5 Reactivity of Lead-Calcium-Tin Alloys
Precise control of the Ca content is required to control the grain structure, corrosion resistance, and mechanical proper-
85
ties of P b - C a - S n alloys. Ca reacts readily with air and other elements such as Sb, As, S to produce oxides or intermetallic compounds which are either lost from the melt (float) or remain suspended. Ca can be lost rapidly from lead (Young and Barclay, 1973). P b - C a alloys can be protected against Ca loss by the addition of Al to the alloy (Prengaman, 1984 a). Al provides a protective oxide skin on molten Pb-Ca-(Sn) alloys. Alloys without Al lose Ca rapidly while alloys containing about 0.0125% or more Al exhibit negligible losses. Lower levels are less protective (Prengaman, 1992). Ca may be added by use of a 73% Ca/27% Al master alloy. With this method, both Ca and Al additions can be controlled. Temperatures must be controlled to retain the Al in solution in the alloys (Prengaman, 1992 b).
2.4 Lead-Tin Alloys The major use for Pb-Sn alloys is for solders and low melting-point alloys for sealing and joining other metals. The P b Sn phase diagram is shown in Figure 2-6 (Hansen, 1958). The Pb-Sn system exhibits a eutectic point at 61.9% Sn at 183 °C. The solid solubility of Sn at 183 °C is 19% decreasing to 1.9% at 25 °C. The decrease in solubility with temperature permits Pb-Sn alloys to be hardenable via precipitation of Sn (Tu, 1972; Tu and Turnbull, 1967, 1969). In Pb-Sn alloys with 2%-20% Sn, the Sn precipitates form solid solution via a discontinuous precipitation reaction. The grain boundaries as well as the cell boundaries resulting from Sn segregation during coring facilitates the nucleation and multiplication of the lamellar precipitates. The precipitates occur at the (010) and
86
2 Lead Alloys
10
Pb
20 30
10
20
40
Tin, at. % 60 70
50
30
40
SO
90
50 Tin, wt %
Figure 2-6. Lead-tin phase diagram (Hansen, 1958).
(001) faces of the lead crystals. The activation energy for diffusion of tin in leads is 26kcal/mol (Seith and Laird, 1932) and thus the diffusion is too slow at 25 °C for diffusion-controlled precipitation, and discontinuous precipitation becomes the dominant mode of precipitation.
Solders with 35-50% Sn have a wide plastic range and are used for general purpose soldering. Sometimes, 0.5-2.0% Sb is added to improve mechanical properties. Sb contents of 0.2-0.5% prevent brittle phase transformations at low temperature. 2.4.2 Corrosion Protection
2.4.1 Solders
The major use of Pb-Sn alloys is as solders for sealing and joining metals. Solders range in compositions from 20 to 98% Pb, the remainder being tin. Pb-Sn (2%) solder is used to seal the side seams of tinned steel cans. Additions of 0.5% Ag to the alloy significantly improves the creep strength. Eutectic Pb-Sn solder with 63% Sn is used for electronic soldering, particularly for printed circuit boards. Pb alloys with 15-30% Sn are used to seal automobile and other types of heat exchangers.
Pb-Sn alloys are used as corrosion-resistant coating on steel and copper. Terne steel is steel sheet coated with a lead alloy containing 15-20% Sn. This alloy contains sufficient Sn to alloy with the surface of the steel. A similar coating, containing 4% Sn, is applied to copper sheet and is used primarily for building flashings. Other Pb-Sn alloys, usually with 50% Sn, are applied as coatings to steel and copper electronic components for corrosion protection, appearance, and ease of soldering.
2.7 References
2.5 Other Lead Alloys Lead has been alloyed with Te, Cu, Cd, Sr? Bi, Ag, Na, Ba, Li, Mg, In, Tl, Zn? Co, and other elements. Only a few of these alloys have found commercial application as specialty metals. In most cases, the properties and performance of lead alloyed with the above elements can be duplicated with Pb-Sb, Pb-Ca, and Pb-Sn alloys combined with minor alloying elements. Thus, many of the lead alloys of the past have been replaced.
2.6 Concluding Remarks Most of the lead alloys used in commerce are cast. Pb-Sb alloys form the bulk of these castings. Because they are strengthened by the as-cast eutectic, Pb-Sb offer superior handling to other lead alloys. At high levels (> 5%), the alloys resist recrystallization despite the low melting point because of the presence of large amounts of second-phase Sb eutectic. The presence of this eutectic phase in finely divided form is responsible for the high creep strength, yield strength, and hardness of these alloys. Unfortunately, most of the applications for cast Pb-Sb alloys apart from lead-acid batteries are decreasing rapidly. Cast Pb-Sb alloys are rapidly losing favor as grid materials in lead-acid batteries because antimony causes gassing during recharge. The market is moving rapidly to maintenance-free SLI batteries, and sealed valve-regulated standby power, telecom, motive power, and portable power batteries. These batteries require alloys for the current carrying and structural members without Sb such as P b - C a alloys. Unfortunately, the yield strength and creep resistance of the Pb-Ca alloys in cast form are
87
significantly inferior to those of cast Pb-Sb alloys. The wrought P b - C a - S n alloy offer significantly higher mechanical properties and offer the opportunity for continuous processing from coils of wrought strip. The addition of other alloying elements to the wrought lead calcium in alloys offers the opportunity for even higher mechanical properties without sacrificing the "maintenance free" aspect of batteries produced from these alloys. Lead-alloy research is now focusing on the mechanical properties and corrosion resistance of cast and wrought P b - C a Sn-(x) alloys. If the research is successful, superior, lighter weight lead-acid batteries will result. Lead is rapidly becoming a one-market commodity (lead-acid batteries). If lead is to survive the competition from lighter, higher power batteries, the mechanical properties of the lead alloys for batteries must be improved.
2.7 References Aballe, M., Redigor, I T., Sistiaga, X M., Torralba, M. (1972), Z. Metallkde. 63, 564. Adeva, P., Caruana, G., Aballe, M., Torralba, M. (1980), Pb80. London: Lead Development Assn., p. 48. ILZSG (1994), Annual Statistics, London: Internal Lead/Zinc Study Group. ASTM (1993), White Metal Bearing Alloys, ASTMB23-83. Philadelphia, PA: ASTM, Vol. 02.04, Sect. 2,9. Bagshaw, N. E. (1968), Lead 68. London: Lead Development Assn., p. 20. Bagshaw, N. E. (1989), /. Power Sources 12,113. Bagshaw, N. E. (1991), J. Power Sources 33, 3. Baker, I. (1993), Acta Metall. Mater. 41, 2633. Barclay, J. B., Ling, F. W, Zeman, R. W. (1980), Lead 80, London: Lead Development Assn., p. 27. Battery Council International (1993), Report of Statistics Committee, Annu. Mtg., Phoenix. Chicago: BCI, p. 15. Bluth, M., Hanemann, H. (1937), Z. Metallkde. 29, 44.
88
2 Lead Alloys
Boehnke, R. W. (1991), Recycling Lead and Zinc, Rome, 1991. London: ILZSG, p. 1. Borchers, H., Assmann, H. (1978), Z. Metallkde. 69, 43. Borchers, H., Assmann, H. (1979), Metall 33, 936. Borchers, H., Kainz, M. (1965), Metall 19, 103. Borchers, H., Reuleaux, M. (1980), Metall 32, 811. Borchers, H., Scharfenberger, W. (1966), Metall 20, 811. Borchers, H., Scharfenberger, W., Henkel, S. (1975), Z. Metallkde. 66,111. Bouirden, L., Hilger, J. P., Hertz, J. (1991), J. Power Sources 33, 27. Burbank, J. B. (1964), /. Electrochem. Soc. 11, 1112. Caillerie, J. L., Hertz, X, Boulahrouf, A., Dirand, M., Hilger, J. P. (1986), Lead 86, London: Lead Development Assn., p. 57. Caldwell, T. W, Sokolov, U. S. (1976), /. Electrochem. Soc. 123, 972. Cole, E. R., Lee, A. Y. (1983), /. Met. 8, 42. Dawson, J. L., Gillibrand, M. T., Wilkinson, J. (1970), at 7th Power Sources Symp., Brighton, U.K. Dean, R. S., Ryjord (1930), Met. Alloys 1, 410. Dean, R. S., Zickrick, L., Nix, F. C. (1926), Trans. AIME 73, 505. Eby, D. J. (1990), in: Lead-Zinc 90. Warrendale, PA: TMS, p. 825. Forest, H., Wilson, X D. (1990), in: Lead-Zinc 90. Warrendale, PA: TMS, p. 971. Giess, H. K. (1984), in: Advances in Lead-Acid Batteries: Pavlov, D., Bullock, K. (Eds.). Pennington, NJ: The Electrochemical Society, 84-14, 241. Gower, R. A. (1990), in: Metals Handbook, 10th ed., Metals Park, OH: ASM, p. 1233. Greenwood, J. N., Cole, J. H. (1949), Metallurgica 39, 241. Hansen, M. (1958), Constitution of Binary Alloys, New York: McGraw-Hill, p. 405, 1101, 1107. Haring, H. E., Thomas, U. B. (1935), Trans. Electrochem. Soc. 68, 293. Hertz, X, Fornasieri, C , Hilger, J. P., Notin, M. (1993), in: Proc. Labat '93, Varna, Bulgaria. Sofia, Bulgaria: Central Laboratory of Electrochemical Power Sources, p. 42. Heubner, U., Sandig, H. (1971), Lead 71. London: Lead Development Assn., p. 29. Heubner, U., Ueberschaer, A. (1974), Lead 74. London: Lead Development Assn., p. 59. Heubner, U., Mueller, I., Ueberschaer, A. (1975), Z. Metallkde. 66, 73. Kallup, B. E. (1986), Lead86. London: Lead Development Assn., p. 68. Kallup, B. E., Berndt, D. (1984), in: Advances in Lead-Acid Batteries: Pavlov, D., Bullock, K. (Eds.). Pennington, NJ: The Electrochemical Society, 84-14, 241. Kostic, N. M. (1982) at AIME Annu. Mtg., Dallas, TX, Febr. 1982. Mao, G. W, Larson, X G. (1968), Metallurgica 3, 236.
Mao, G. W, Larson, X G., Rao, P. (1969), J. Inst. Met. 97, 343. Myers, M., Van Handle, H., Dimartini, C. (1974), /. Electrochem. Soc. 121, 11. Niragu, X (1978), The Biogeochemistry of Lead. New York: Elsevier. Olper, I. M. (1991), in: Recycling Lead and Zinc, Rome, 1991. London: ILZSG, p. 79. Olper, I. M., Asano, B. (1989), in: Primary and Secondary Lead Processing. New York: Pergamon, p. 119. OSHA (1978), Occupational Exposure to Lead, Appendix A29CFR 1910.1025, U.S. Federal Register, Washington, D.C. Pavlov, D. (1991), J. Power Sources 33, 221. Pavlov, D. (1992), in: Proc. 6th ILZRO, Lead Acid Battery Seminar, Research Triangle Park, NC: ILZRO, p. 169. Perkins, X, Pokorny, X L., Coyle, M. T. (1976), Report No. NPS-69PS76101, Naval Research Laboratory, Washington D.C. Piercy, R. C. (1976), at Fall Mtg. Electrochem. Soc, Las Vegas, October 1976. Pierson, X, Weinlein, C , Wright, C. (1975), in: Power Sources 5. New York: Academic Press, p. 97. Pike, K. N. (1990), Lead-Zinc 90. Warrendale, PA: TMS, p. 955. Prengaman, R. D. (1976), at Fall Mtg. Electrochem. Soc, Las Vegas, NV, October 1976. Prengaman, R. D. (1977 a), Proc. IBM A, Largo, FL, p. 121. Prengaman, R. D. (1977 b), at Fall Mtg. Electrochem. Soc, Atlanta, GA, October 1977. Prengaman, R. D. (1980), in: Lead80. London: Lead Development Assn., p. 34. Prengaman, R. D. (1983), in: Lead83. London: Lead Development Assn., p. 69. Prengaman, R. D. (1984a), in: Advances in LeadAcid Batteries. Pavlov, D., Bullock, K. (Eds.). Pennington, NJ: The Electrochemical Society, 84-14, p. 201. Prengaman, R. D. (1984 b), in: Anodes for Electrowinning: James, S. (Ed.). Warrendale, PA: TMS, p. 49. Prengaman, R. D. (1987), CU87, Cooper, W, Lagos, G. E., Ugarte, G. (Eds.). Santiago: University of Chile, Vol. 3, p. 387. Prengaman, R. D. (1989), in: Proc. 3rd Int. LeadAcid Battery Seminar. Research Triangle Park, NC: ILZRO, p. 1. Prengaman, R. D. (1991 a), in: Recycling Lead and Zinc, Rome. London: ILZSG, p. 437. Prengaman, R. D. (1991 b), /. Power Sources 33, 13. Prengaman, R. D. (1992 a), Batteries Int. 13, 24. Prengaman, R. D. (1992 b), Batteries Int. 10, 52. Prengaman, R. D. (1993), J. Power Sources 42, 25. Prengaman, R. D., McDonald, H. (1990), in: Lead and Zinc 90: Mackey, T, Prengaman, R. (Eds.). Warrendale, PA: TMS, p. 1045.
2.7 References
Rose, M. V., Young, J. A. (1974), Lead 74. London: Lead Development Assn., p. 37. Ruetschi, P., Cahan, B. D. (1975), /. Electrochem. Soc. 104, 406. Rutter, M., Jones, R. R. (1983), Lead vs. Health Sources and Effects of Low Level Lead Exposure. New York: Wiley. Scharfenberger, W, Henkel, S. (1973), Z. Metallkde. 64, 478. Schuhmacher, E. E., Bouton, G. M. (1930), Met. Alloys 1, 405. Schuhmacher, E. E., Phipps, G. S. (1935), Trans. Electrochem. Soc. 68, 309. Seith, W, Laird, I G. (1932), Z. Metallkde. 24, 193. Serracane, C. (1991), in: Recycling Lead and Zinc, Rome. London: ILZSG, p. 91. Simon, A. C. (1965), Batteries 2. New York: Pergamon. Smith, J. W. (1979), in: Metals Handbook, Vol. 2, 9th ed. Metals Park, OH: ASM, p. 511. Taiton, U. C. (1928), Eng. Min. J. 1, 9. Townsend, A. B. (1960), AEC Research and Development Report, Y-1307 Metallurgy and Ceramics, Union Carbide Contract W7405 Eng., p. 26. Tu, K. N. (1972), Met. Trans. 3, 2769. Tu, K. N., Turnbull, D. (1967), Ada Metall. 15,1317. Tu, K. N., Turnbull, D. (1969), Acta Metall. 17,1263. U.S. Bureau of Mines (1964, 1993), Mineral Industry Survey, U.S. Dept. of Interior, Washington, DC. Vinal, G. W (1955), Storage Batteries, 4th ed. New York: Wiley.
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Wade, E. J. (1902), Secondary Batteries. London: The Electrician Printing & Publishing Co. White, J. W R., Rogers, R. (1974), Lead 74. London: Lead Development Assn., p. 67. Willingantz, E. (1972), Telephony 12, 17. Worchester, A. W, O'Reilly, J. T. (1990), in: Metals Handbook, Vol.2, 10th ed. Metals Park, OH: ASM, p. 543. Young, J. A., Barclay, J. B. (1973), 85th Annu. Mtg., BCI Conf. Proc. Chicago: BCI. Zintl, E., Neumayr, S. (1933), Z. Electrochem. 39, 86.
General Reading Hoffmann, W. (1970), Lead and Lead Alloys. New York: Springer. Kuhn, A. T. (1979), The Electrochemistry of Lead. New York: Academic Press. Prengaman, R. D. (1990), "Recycling of Lead", in: Metals Handbook, Vol.2; 10th ed. Metals Park, OH: ASM, p. 1221. Smith, X W (1979), "Corrosion Resistance of Lead", in: Metals Handbook, Vol. 2, 9th ed. Metals Park, OH: ASM, p. 511. Vinal, G. W. (1955), Storage Batteries, 4th ed. New York: Wiley. Worchester, A. W, O'Reilly, J. T. (1990), "Lead and Lead Alloys", in: Metals Handbook, Vol. 2, 10th ed. Metals Park, OH: ASM, p. 543.
3 Zinc Dimitri Coutsouradis and Gisele Walmag Centre de Recherches Metallurgiques (C.R.M.), Liege, Belgium
List of 3.1 3.1.1 3.1.2 3.1.2.1 3.1.2.2 3.1.2.3 3.2 3.2.1 3.2.2 3.2.2.1 3.2.2.2 3.2.2.3 3.2.2.4 3.3 3.3.1 3.3.1.1 3.3.1.2 3.3.1.3 3.3.1.4 3.3.2 3.3.2.1 3.3.2.2 3.4 3.4.1 3.4.2 3.4.3 3.4.3.1 3.4.3.2 3.4.4 3.5 3.6 3.7 3.8
Symbols and Abbreviations 92 Introduction 93 Main Areas of Applications 93 Technical Requirements for the Applications 93 Physical Properties 94 Electrochemical and Chemical Properties 94 Alloying Behavior 95 Galvanizing Alloys 96 Protection of Steel by Means of Zinc Coatings 96 Hot Dip Galvanizing Processes 96 Hot Dip Galvanizing Process After Manufacturing (Batch Galvanizing) .. 96 Continuous Hot Dip Galvanizing Process 97 Zinc Alloys for Galvanizing After Manufacturing 98 Zinc and Zinc Alloy Baths for Continuous Hot Dip Galvanizing 99 Casting Alloys 100 Pressure Die-Casting Alloys 100 The Pressure Die-Casting Process 100 Composition, Microstructures, Properties 101 Impurities 101 Mechanical Properties 102 Gravity Casting Alloys 102 Physical and Mechanical Properties 102 Wear Properties 103 Wrought Alloys 103 Plastic Deformation, Recovery and Recrystallization of Zinc 103 Alloys 105 Casting and Rolling Processes 106 Casting 106 Rolling 106 Properties 107 Recycling 108 Concluding Remarks 108 Acknowledgement 109 References 109
Materials Science and Technology Copyright © WILEY-VCH Verlag GmbH & Co KGaA. All rights reserved.
92
3 Zinc
List of Symbols and Abbreviations a c m n Tm
lattice parameter lattice parameter strain rate sensitivity strain sensitivity melting temperature of zinc in kelvin
a F 8 s £ r|
f.c.c. Al-rich phase intermetallic compound intermetallic compound intermetallic compound intermetallic compound pure zinc phase
DDR DRR DRV DRX EAF f.c.c. h.c.p. IACS RT SEM SFE TEM
discontinuous dynamic recrystallization dynamic rotation recrystallization dynamic recovery dynamic recrystallization electric arc furnace face centered cubic hexagonal close packed international annealed copper standard room temperature scanning electron microscopy stacking fault energy transmission electron microscopy
Fe3Zn10 FeZn10 CuZn4 FeZn13
3.1 Introduction
3.1 Introduction The world consumption of zinc was 7 Mt in 1990. In comparison, the consumption of aluminum, copper and lead was 17.9, 10.8 and 9 Mt, respectively (Metallstatistik, 1991). During the period 19801990, zinc consumption experienced a growth of 2.2% while the growth rate of the other nonferrous metals aluminum, copper and lead was 2.1, 1.6 and 1.5%, respectively. Copper and lead behaved during the 1980s like more mature metals (Masson, 1991). 3.1.1 Main Areas of Applications
It is usual to describe the applications of zinc in terms of products. The breakdown for the consumption of the main products in some industrialized countries in 1990 is summarized in Table 3-1. Zinc consumed for galvanizing purposes represents by far the largest application for zinc: about 50% of the total consumption. Galvanizing after fabrication and continuous galvanizing share this field about equally. Casting alloys constitute the second biggest application area for zinc, account-
Table 3-1. Main product areas for applications of zinc in six industrial countries in 1990 (FRG, F, UK, I, J, USA) (Metallstatistik 1980-1990, 1991). Product area
Tonnage
643416 Brass 1556 396 Total galvanizing3 192900 Semi products (sheet,...) 228 693 Zinc oxide and compounds Casting alloys 591 908 156831 Miscellaneous 3 370144 Total a
Percentage 19.09 46.18 5.72 6.79 17.56 4.65 100.00
Sheet, after fabrication, pipes, wire, the first two being by far the most important.
93
ing for about 17.5% of the zinc consumption. This comprises essentially zinc alloys for pressure die-casting, although in recent years new alloys have found applications as gravity castings. Wrought alloys account for almost 6% of the zinc consumption and comprise mainly sheet for gutters, roofing and other building applications, which are quite widespread in some countries, especially in Europe. Other applications for wrought alloys comprise zinc sheet for battery cells and, in some countries, sheet for coins. An interesting application of sheet, although of small tonnage, is the fabrication of fuses especially for the automotive industry. Wrought alloys produced by deformation processes other than rolling have been studied and alloys specific for this application have been developed but have not achieved any industrical importance. Zinc wire and zinc alloy wire are extensively used for corrosion protection of steel structures by metallizing. Zinc powder is an important chemical reactant used in great amounts by the zinc industry and in smaller amounts by other industries. Zinc powder is also used in paints. Among the zinc compounds, zinc oxide finds wide usage in the rubber industry, and to a lesser extent in paints and other industries. In this chapter we shall deal mainly with galvanizing, casting alloys and wrought alloys. Applications of zinc as an alloying element in copper (brass) or in aluminum alloys is dealt with in other chapter of this volume. Zinc in brass represents 19% of the uses of the metal. 3.1.2 Technical Requirements for the Applications
Applications of zinc, whatever the area, are justified on the basis of technical char-
94
3 Zinc
acteristics of zinc and, of course, economic reasons. The technical reasons for the use of zinc, zinc alloys or zinc compounds are based on the properties of the respective product. The properties of zinc which are relevant as regards applications may be classified mainly as physical properties, and chemical, electrochemical and alloying behavior.
3.1.2.1 Physical Properties
Table 3-2 lists the physical properties of the element zinc. Some of these properties are relevant to the applications of zinc as discussed hereafter. The low melting temperature of zinc (420 °C), together with the low heat of fusion and the modest heat capacity, accounts for the success of zinc as a casting alloy in the hot chamber die-casting process (see below). The low melting point of zinc is also an asset for the main application of zinc in hot dip galvanizing. On the other hand, the low melting temperature is an inconvenience because it limits the mechanical properties at elevated temperatures. The coefficient of thermal expansion and its anisotropy has a bearing in sheet applications for roofing and gutters. Suitable electrical resistivity and melting point are the properties required in an application such as fuses. The fairly high thermal conductivity is of interest in the fabrication of parts, castings for example, subjected to heat. The formability of zinc is typical of that of a hexagonal metal with a high c/a ratio. This applies especially to low-temperature deformation. At high temperatures the formability of zinc is excellent in all cases for which it is needed, i.e., the hot rolling of zinc sheet.
Table 3-2. Physical properties of zinc (ASM, 1990). Unit
Property
°C °C
close packed hexagonal a=0.26648 c = 0.4947 c/a = 1.856 65.38 7.133 419.58 906
%
4.7
kJkg^K"1
0.417
Crystal structure Lattice parameters
Atomic weight Density at 25 °C Melting point Boiling point Volume contraction on solidification Mean specific heat (100-200 °C) Thermal conductivity at 25 °C Linear thermal expansion (single crystal 0-100°C)
Value
nm
g/cm3
¥ K
1
Linear thermal expanK" 1 sion (polycrystalline solid 0 - 250 °C) Latent heat of fusion kJ/kg Latent heat of vaporization MJ/kg Electrical conductivity %IACS Electrical resistivity (polycrystalline solid at20°C) nQm Modulus of elasticity at 20 °C (no real elastic behavior) GPa
113 along a axis: 15xlO~ 6 along c axis: 61.5 x l O " 6 39.7 x l O " 6 (depends on texture) 100.9 1.782 28.27
59.16
70-140
3.1.2.2 Electrochemical and Chemical Properties
The position of zinc in the electrochemical potential series accounts for the sacrificial protection of steel when the latter is protected by means of a zinc coating (see Chap. 11 of Vol. 15 of this Series). In spite of its high electronegative level, zinc has an excellent corrosion resistance due to the fact that the corrosion products formed in the atmosphere are fairly insolu-
95
3.1 Introduction
ble and thus provide protection from further attack (Christman et al., 1986). These properties are at the basis of all applications concerning the protection of steel and sheet for the building industry and also an important asset for the development of the casting alloys.
3.1.2.3 Alloying Behavior
Because of the crystal structure and the electronic configuration of zinc, the solubility of most elements in solid-phase zinc is very low. Strengthening by solid solution is therefore also limited. No element is known to form an interstitial solid solution. On the other hand there is no allotropic transformation in zinc and the derived strengthening mechanisms (e.g.
martensitic transformation) are not available. With the exception of the Zn-Al system, there are no other soluble elements to provide coherent strengthening precipitates as is the case in aluminum or nickel alloys. The strengthening of zinc alloys relies mainly on solid-solution, grain size and interphase strengthening. The binary phase diagrams (ASM, 1992) of Zn-Al (Fig. 3-1) and Zn-Cu are important considering applications of Zn-Al alloys as casting and coatings and the use of Zn-Cu alloys in sheet and casting applications. In many systems, zinc presents eutectic reactions with limited solubility in the solid state. Among these systems the Fe-Zn system (Fig. 3-2) is of major relevance in hot dip galvanizing alloys whereas the Zn-Ti system is important for sheet alloys.
Atomic Percent Zinc 10
800
20
30
40
50
60
70
90 100
418.58'C
10
Al
20
40
50
60
Weight Percent Zinc
Figure 3-1. The Zn-Al phase diagram (ASM, 1992).
70
80
90
100
Zn
96
3 Zinc
Atomic Percent Zinc 30
10
40
50
70
60
80
90
100
1600-
419.58°C 30010
20
Fe
30
40
50
60
Weight Percent Zinc
70
80
90
100
Zn
Figure 3-2. The Zn-Fe phase diagram (ASM, 1992).
3.2 Galvanizing Alloys 3.2.1 Protection of Steel by Means of Zinc Coatings In many applications, especially in the construction, transportation and agricultural areas, steel does not offer sufficient corrosion resistance and has therefore to be protected. Zinc is a very successful coating for steel for several reasons: • It provides sacrificial protection. • The corrosion resistance of zinc is much higher than that of iron. Indeed, data provided by long time exposure in many sites show that the corrosion rate of zinc is up to 350 times lower than that of iron. Usually one considers that the corrosion rate of zinc is at least 35 times
lower than that of iron. Corrosion rates (in jim/year) of zinc - or pure zinc coatings - are typically as follows for different environments: rural, 1.5; urban, 4 - 5 ; industrial, 5-10; marine, 2 (Christman etal, 1986). • Several methods are available for coating steel with zinc: hot dip, electrolytic coating, vapor deposition, metallizing, mechanical coating. We shall consider here only the hot dip coatings. 3.2.2 Hot Dip Galvanizing Processes 3.2.2.1 Hot Dip Galvanizing Process After Manufacturing (Batch Galvanizing) The parts are first attached to appropriate supports and then, typically, subjected successively to the following treatments:
3.2 Galvanizing Alloys
Layer
97
Hardness DPN 70
179
81 $1K
244
r Fe 16.1 um
degreasing, pickling, fluxing, immersion in the molten zinc (Porter, 1991). The first steps aim at cleaning the surface to insure optimum reactivity conditions between steel and zinc. During immersion, the reaction between zinc and iron occurs after the steel surface has been wetted by the liquid zinc. The reaction rate is parabolic below 490 °C, becomes linear between 490 and 520 °C and then becomes parabolic again at temperatures above about 530 °C. The differences in the reaction rates as a function of temperature are due to changes in the nature of the intermetallic compounds in the Fe-Zn system. The Fe-Zn reaction is described by Horstmann (1978). At normal galvanizing temperatures (460 °C), the zinc coating layer consists of the following intermetallic compounds in order from the surface of the steel: F, 8 and £, covered by a layer of the r| zinc phase. Figure 3-3 shows the microstructure of a typical coating (European Zinc Producers, 1981). For the hot dip coating produced after manufacture, the thickness of the zinc coating layer ranges from about 40 Jim for small parts immersed for a short time to more than 80 \im for heavier parts immersed for several minutes.
159
Figure 3-3. Microstructure of a hot dip galvanized coating; SEM. DPN: diamond pyramid number, which is equivalent to Vickers hardness (European Zinc Producers, 1981).
A major feature of the batch galvanizing process is the effect of steel composition (Si and P contents) on the rate of the Fe-Zn reaction. Figure 3-4 shows the so-called Sandelin effect, which leads to excessive zinc coating thicknesses for Si contents around 0.1 Vo1 (Sandelin, 1940), and that the same effect is actually a function of the parameter Si + 2.5P (Leroy etal, 1976). The excessive reactivity results in thick coatings which are sometimes also very brittle. The role of silicon has been the object of numerous studies. In Sandelin-type steels, it has been shown that the occurrence of a high steel reactivity is due to the precipitation in the coating of Fe-Si particles (Pelerin et al., 1985). 3.2.2.2 Continuous Hot Dip Galvanizing Process The main steps of the continuous galvanizing process are heating, annealing, cooling, immersion in the zinc bath under a reducing (N 2 -H 2 ) atmosphere, removal from the bath and gas wiping to remove excess zinc. In several plants, the sheet is, after wiping, reheated by induction or in gas-fired 1 All compositions, in the text and in the tables, except where otherwise stated, are cited in wt.%.
98
3 Zinc
400 350 300
J-250 CO
S 200 c
I 150 i—
Figure 3-4. Plot of coating thickness vs. Si + 2.5 P content showing effect of Si and P on the Fe-Zn reaction (Leroy et al., 1976).
100 50 0
50
100
150
200
250
300
Si + 2.5P(10"
furnaces in order to transform free zinc to Fe-Zn intermetallics (galvannealing). The Fe-Zn coating obtained in this way is highly appreciated because of its excellent spot weldability and paintability. In continuous sheet galvanizing, the immersion time in the zinc bath is short (a few seconds), the thickness of the intermetallics is therefore low and the thickness of the coating is also low (generally below 20 jum). The gas wiping controls the thickness of the r| phase (solid zinc) in the coating. The occurrence of thick intermetallic layers is, in general, associated with brittleness of the coating. Because most of the coated sheet is subject to deformation during fabrication, it is important to ensure a good ductility of the coating, which implies that the intermetallic layer must be as thin as possible. Apart from bath composition (see below), temperature and immersion time are the main controlling parameters for the thickness of the intermetallic layer. 3.2.2.3 Zinc Alloys for Galvanizing After Manufacturing
The Z n - l P b alloy is the widely used alloy for batch galvanizing. Lead is insoluble in solid zinc. During solidification, liquid lead particles are enclosed in the solid
zinc, the particles (spherical in form) solidifying when the temperature decreses below 318 °C. In the presence of lead in amounts above about 0.1%, the solidification of the coating results in the formation of a typical spangle. The spangle consists of a number of surface grains, some shiny and others frosty. Spangle formation is a result of the dendritic mode of solidification exhibited by alloys with solutes with limited solubility in zinc (Krepski, 1980) and is regarded as most desirable by many users since this aspect characterizes the zinc coating. From the point of view of practical experience, it is reported that lead decreases the surface tension of the bath (Krepski, 1980) and thus that it improves wetting of the steel part. A layer of lead at the bottom of the galvanizing kettles is also helpful for the removal of dross (Fe-Zn intermetallics). Aluminum is added to the zinc bath in small amounts (0.007% maximum) in order to improve the brightness of the coating. In higher amounts, aluminum tends to react with the fluxes, which in turn gives rise to surface defects. No other intentional addition elements are used in zinc baths. However, depending on the raw materials used, other elements
3.2 Galvanizing Alloys
such as tin, cadmium and antimony may be present in small amounts (0.001 -0.01 %). Tin is reported to enhance spangle and to increase the brightness. Antimony also enhances spangle while cadmium has no significant effect, at least when present as a residual element. The Polygalva alloy (Dreulle, 1979; Pelerin, 1983) (0.04%Al-0.007%Mg-~ 0.03%Sn-0.3%Pb) was developed specifically for the purpose of resolving the problem of the Sandelin peak. Aluminum is present in order to reduce the Fe-Zn reaction rate. Tin and magnesium are added in order to improve the wettability and therefore the appearance of the coating. In the Technigalva bath (0.04%Ni-Pb), the active element is nickel. Nickel reduces the coating thickness at the Sandelin peak, has no effect on the reactivity of high-silicon steels and reduces the layer thickness of the silicon-free steels (Sokolowski, 1988). 3.2.2.4 Zinc and Zinc Alloy Baths for Continuous Hot Dip Galvanizing
The Zn-Al system provides all of the alloys used in hot dip galvanizing. Lowaluminum alloys with aluminum contents ranging from 0.15 to 0.25% without any other addition (iron at the saturation level is however inevitably present) are used for spangle-free coatings as required for applications in the automotive industry (Meuthen, 1989). Similar baths with additions of lead (0.1 % and higher) are used for the production of coatings with a spangle as still required by many users. The role of aluminum in these coatings is essential. At low contents, the aluminum inhibits growth of intermetallics and thus ensures a coating as ductile as possible. The actual mechanism of the inhibition is rather complex. It depends on the formation of ternary Fe-Zn-Al phases during the reaction. It has been observed that at
99
the very initial stages of the reaction an Fe 2 Al 5 intermetallic phase forms at the Zn-Fe interface, which slows down the diffusion of species for the formation of F e Zn intermetallic layers. The thin (or not so thin) intermetallic layer consists of the known phases F, 5 and £. The special features of the Zn(Al)-Fe reactions in continuous galvanizing depend on the conditions, temperature, duration of contact, etc. and also on the composition of the steel (Hisamatsu, 1989). Another function of aluminum is to act as a "reducing" agent for some of the oxides still remaining on the steel surface. Indeed, it has been shown that the wettability of the bath increases with the aluminum content up to a certain content. At higher aluminum contents (above about 1 %), the wettability of the bath decreases due to surface tension and fluidity effects (Renaux and Silvestre, 1988). A high aluminum content is not desirable in terms of kinetics when a galvannealing reaction follows the immersion phase. In such a case, if the aluminum content in the bath is reduced to 0.1% or less the galvannealing reaction is easier but at the expense of the soundness of the coating (powdering). Galvannealing consists in the transformation by diffusion processes of the zinc layer to 8, £ and F intermetallics, the relative proportion and microstructure of which are responsible for the properties of the coating. Considerable work has been devoted to the study of the kinetics and to the control of the galvannealing reaction (Lamberigts et al., 1992). Higher aluminum contents are used mainly in order to enhance the corrosion resistance of the coating. Such higher contents are first the 5% eutectic level (Galfan and Superzinc) and then the 34.5Zn-1.5Si-Al bath (Galvalume and other names).
100
3 Zinc
The Galfan alloy, introduced first in 1982, contains along with 5%A1 also small additions of mischmetal intended to improve the wettability of the alloy (Coutsouradis etal., 1984; Coutsouradis and Ayoub, 1989). Superzinc contains magnesium additions claimed to improve the paintability of the Zn-4%A1 alloy. The Galfan coating is characterized by the virtual absence of any extensive intermetallic layer, which, together with the deformation characteristics of the alloy, provides a most ductile coating. The alloy has proved particularly suitable for subsequent coil coating (Schwarz et al, 1992). The 55%A1-Zn coating, introduced in 1964, also contains 1.5% Si intended to inhibit the Fe-Al reaction. The main feature of this coating is its excellent corrosion resistance, arising from the high aluminum content, combined with sacrificial protection, owing to the presence of zinc. This coating has gained wide acceptance especially in building applications (Kriner and Niederstein, 1991).
3.3 Casting Alloys Zinc alloys are extensively used in castings and can be divided in two families
depending on their composition and the casting process. The first group of alloys comprises the Zn-4%A1 alloys first used 60 years ago and cast in hot chamber pressure die-casting machines. The second group is composed of more recently developed alloys with aluminum contents of 8, 12 and 27% which can also be gravity cast. Table 3-3 lists the nominal compositions of these alloys. 3.3.1 Pressure Die-Casting Alloys 3.3.1.1 The Pressure Die-Casting Process The hot chamber pressure die-casting process is characterized by the fact that the piston is permanently immersed in molten metal. During recent decades the process has undergone considerable evolution in terms of technology and modeling of the die design (Porter, 1991; Ogier, 1993; Spittle and Brown, 1993). The thin wall die-casting technology allowed the zinc alloys to successfully resist the inroad of plastics because of the reduction of weight achieved (Radtke, 1972). The use of an alloy in the hot chamber die-casting machine requires a low corrosion rate of tool steels in the liquid metal.
Table 3-3. Nominal composition of casting alloys (%) (Porter, 1991). Alloy 2 3 5 7 ZA-8 ZA-12 ZA-27 ILZRO 14 ILZRO 16
Al
3.5-4.3 3.5-4.3 3.5-4.3 3.5-4.3 8 -8.8 10.5-11.5 25 -28 0.02
Cu 2.5-3.5 0.25 max. 0.5-1.25 < 0.25/0.012 Ni 0.8-1.3 0.5-1.2 2.0-2.5 1.25 1.25
Mg
0.02 -0.06 0.02 -0.06 0.03 -0.08 0.005 0.015-0.03 0.015-0.03 0.01 -0.02 0.13 Ti 0.27Ti/0.2Cr
Density in g cm ~ 3 6.6 6.6 6.7 6.7 6.3 6.0 5.0 7.1 7.1
3.3 Casting Alloys
In most zinc alloys in which aluminum constitutes between 4 and 8%, the presence of this element, associated with a low melting temperature, results in a very slow corrosion rate of the tools. Alloys ZA-12 and ZA-27, which have a higher aluminum content, have a fairly high melting temperature, which therefore induces corrosion of the tools. On the other hand, the alloys ILZRO 14 and 16, because they contain no aluminum, also attack the tools and therefore have to be cast in cold chamber machines in which the molten metal is maintained in a separated vessel and then transferred to the piston for injection. The beneficial effect of aluminum on the corrosion rate of tool steels is due to the formation and the stability of Fe 2 Al 5 or FeAl3 (Galvez-Soza et al., 1984). 3.3.1.2 Composition, Microstructures, Properties
All the zinc casting alloys, except one (ILZRO 14/16), are based on the Zn-Al or Zn-Al-Cu systems. It was discovered early in the development of zinc casting alloys that aluminum alloyed in substantial amounts to zinc can significantly improve the fluidity, refine the cast grain structure and thus enhance the mechanical properties. The Zn-Al equilibrium diagram (Fig. 3-1) shows a binary eutectic at 5% Al with a melting point of 382 °C. Hypoeutectic alloys solidify with a primary zinc-rich r| phase surrounded by eutectic r| + a (Alrich phase), while in hypereutectic alloys a phase is formed first. At 275 °C, a eutectoid decomposition induces the formation of r| and a phases richer in zinc and aluminum, respectively. The microstructure of pressure die-castings is often metastable due to rapid cooling, and the various structural modifica-
101
tions occur within a month or a few years for some of them. For example, aging reactions in Zn-27%A1 alloys involve complex precipitation and transformation phenomena (Lecomte-Beckers etal., 1989). The composition of the hypoeutectic alloys defined around 1940 (Burkhardt, 1940) is the result of a compromise between several properties, mainly strength, ductility, impact resistance, susceptibility to aging, proneness to intergranular corrosion and fluidity. Copper is present in the r| phase in solid solution and restrains solid-state r| decomposition. On cooling, solid solubility of copper in r| decreases and the 8 phase precipitates. Magnesium is an alloying element added to most zinc casting alloys in the range 0.03-0.06% to increase strength, inhibit intergranular corrosion and restrain r\ transformation in Zn-Al alloys. Higher magnesium contents induce lower impact strength and hot tearing due to the presence of a ternary (Zn, Al, Mg) low melting point eutectic. Several addition elements have been studied for their effect on the creep strength of casting alloys: chromium, lithium and manganese (Walmag et al., 1987). Lithium and manganese additions to ZA-8 resulted in the development of the hot chamber pressure die-casting alloy ZATEC (Barnhurst, 1992, 1993), which exhibits an elevated creep strength due to intermetallicphase dispersion hardening. 3.3.1.3 Impurities
Deleterious impurities possibly present in zinc either introduced by alloying elements or picked up from contamination are lead, tin, cadmium, indium and thallium. These elements have a detrimental effect on intergranular corrosion resistance
102
3 Zinc
of alloys that also contain aluminum. The role of the deleterious impurities consists in the precipitation at the grain boundaries of impurity-rich films, resulting in the formation of anodic areas relatively to the matrix (Niessen and Wineguard, 1966). The additions of magnesium and copper decrease the proneness to intergranular corrosion, as mentioned above. 3.3.1.4 Mechanical Properties
Tensile strength increases with aluminum and copper content but decreases with in-40 0
80
240
160
500
320
+ • • o
ZA-27 ZA-12 ZA-8 No. 3 D No. 5
400
300
70
3.3.2 Gravity Casting Alloys
60
3.3.2.1 Physical and Mechanical Properties
50
The hypereutectic Zn-Al alloys were initially developed for gravity casting, i.e., sand casting and permanent mold casting (Gervais et al., 1985). Their composition is identical in both pressure die-casting and gravity casting (Table 3-3). The freezing range of the alloys increases with aluminum content from 29 °C (ZA-8) to 45 °C (ZA-12) to 108 °C (ZA-27). The phases constituting their microstructure are also identical, although their size and shape are different.
40 30
200
20 100 10
_
creasing temperature, leading to a marginal difference at 150°C (Fig. 3-5). A recent study of hypoeutectic alloys and ZA-8 showed that strength is strongly dependent on casting thickness: the thinner the casting the higher the strength, as shown in Fig. 3-6 (Walmag et al., 1991; Goodwin et al., 1993). This effect is observed both at room temperature and at 80 °C. Casting parameters also affect tensile strength. Increasing the velocity with which the metal enters the "die" or decreasing the die temperature allows the strength to be increased but substantially less than the change caused by the casting thickness.
' 160 Temperature, °C
Figure 3-5. Effect of temperature and composition on tensile strength of pressure die-casting alloys (Goodwin and Ponikvar, 1989). 450
400
S. 350
300
250 0,75
1,5
2,25
Thickness(mm) ZA-8
Figure 3-6. Effect of thickness on strength of diecasting alloys (Walmag et al., 1991; Goodwin et al., 1993).
3.4 Wrought Alloys
Consequently, the mechanical properties of ZA-8, ZA-12 and ZA-27 vary considerably with the casting process (Fig. 3-7). For alloy ZA-27, the microstructure and especially the grain size depend on the cooling rate, which thus influences the mechanical properties: strength and ductility increase when the grain size decreases. Alloying elements such as titanium and boron where shown to have a refining effect (Skenazi et al, 1983) and therefore to improve the ductility of sand or permanent-mold-cast ZA-27. The effect of titanium and boron was explained partly by the refinement of the microstructure and partly on the basis of modifications in the eutectoid decomposition (Lamberigts etal., 1985). The ZA alloys are characterized by relatively low densities, which make them of interest for applications where weight is an important parameter. The thermal expansion of ZA alloys is somewhat higher than that of aluminum alloys, copper alloys and cast iron but the thermal and electrical conductivities compare favorably with most aluminum alloys and even exceed those of copper alloys. The ultimate tensile strength of ZA alloys is higher than that of most widely used
103
nonferrous alloys and is equal to or even exceeds that of many cast irons. Moreover, they maintain relatively high strength when the temperature increases. The ultimate tensile strength of ZA-27 at 100 °C remains above that of aluminum, copper and iron gravity casting alloys. Room temperature aging of ZA alloys does not influence their strength, while aging at 95 °C slightly decreases it. 3.3.2.2 Wear Properties
Alloys ZA-12 and ZA-27 sand, die or permanent mold cast offer interesting properties as bearing alloys. The high compression strength and hardness associated with a multiphase microstructure account for the good bearing properties (Barnhurst, 1990 a). However, compared to bronze they are limited by the higher rate of softening with temperature (Kingsbury, 1992).
3.4 Wrought Alloys 3.4.1 Plastic Deformation, Recovery and Recrystallization of Zinc
The plastic deformation and thermally induced softening of zinc and zinc alloys
o 450
ZA-8
ZA-12 ! sand cast
! presssure die-cast
Zk-27
Figure 3-7. Effect of casting process on strength of alloys ZA-8, ZA-12 and ZA-27 (data from Goodwin and Ponikvar, 1989).
104
3 Zinc
are important for the behavior in service of industrial alloys and, of course, for the forming by means of plastic deformation. Volume 6 (Plastic Deformation and Fracture of Materials) and Volume 15 {Processing of Metals and Alloys) of this Series deal with the general and fundamental aspects. In this section, some of the data relevant to zinc will be summarized. The plastic deformation of single crystals of zinc has been extensively studied along with the deformation of other hexagonal metals. As for other hexagonal metals, the deformation of zinc takes place by (Michel, 1979; Wegria, 1985; McQueen and Bourell, 1988; Partridge, 1967; Jaoul, 1965): • gliding along basal planes (0001) <1120>, along prismatic planes (1010)<1120>, along pyramidal planes of the first order {1010} <1120> and along pyramidal planes of the second order {1122}
, • twinning, • kinking. Gliding along basal planes is the predominant deformation mode as expected for an h.c.p. metal with a high c/a ratio (1.856). The small number of gliding planes explains the high tendency of zinc to twinning. The high stacking fault energy (300 erg/cm2; Partridge, 1967) favors gliding along the pyramidal planes of the second order (Dupouy and Poirier, 1968). For zinc polycrystals, deformation takes place predominantly along the basal planes. Other mechanisms can occur depending on stress conditions, temperature, etc.: gliding along pyramidal planes, twinning or intergranular creep. The latter provides a major contribution to deformation at least at temperatures higher than the ductile-brittle fracture transition in zinc, i.e. above 0.33 Tm (-44°C) (Jaoul, 1965; Jacquerie, 1966).
The deformation mechanisms of zinc or zinc alloys with the associated softening have been described in terms of three activation energy levels. For pure zinc, creep tests at temperatures above 0.7-0.8 Tm showed an activation energy of 160 kJ/mol associated with dynamic recrystallization and with the appearance of prismatic slip (McQueen and Bourell, 1988). At lower temperatures of 0.6 Tm to 0.8 Tm, an activation energy of 93 kJ/mol was determined, which compares well with the activation energy for self-diffusion of zinc (97 kJ/mol). The same order of magnitude of the activation energy was found for creep of a polycrystalline Zn-Ag alloy above 0.6 Tm (Davies and Williams, 1975) and by Cosse et al. (1980) in rolled Zn-Cu-Ti alloy and in a die-cast Zn-Al-Cu alloy. The above mentioned studies report also, at temperatures below 0.6 Tm and relatively high strain rates, activation energies of about 54-60 kJ/mol for the alloys studied by Cosse et al. (1980), which corresponds to the activation energy for the diffusion of zinc at the grain boundaries (60kJ/mol; Wajda, 1954). Due to the related low melting temperature and high c/a ratio, zinc recrystallizes at fairly low temperatures. For the pure metal, elimination of deformation dislocations can occur at room temperature. Dynamic recovery (DRV) occurs above 0.25 Tm (-100°C in the case of zinc). Above 0.5 Tm (73 °C), recovery is controlled also by diffusion mechanisms. The softening behavior of zinc is influenced by its stacking fault energy. A high stacking fault energy indicates that dislocations are unable to dissociate and therefore their rearrangement by climbing is made easy (see Sec. 9.3 in Vol. 15 of this Series). Although DRV may be the dominant softening mechanism for pure zinc, dy-
3.4 Wrought Alloys
namic recrystallization by rotation (DRR) was reported in the range 100-300°C (Humphreys, 1981). DRR is due to difficulty of non-basal slip causing lattice rotation in the grain boundary regions. High stacking fault energy metals such as aluminum and zinc would show only DRR while low SFE metals would exhibit discontinuous dynamic recrystallization (DDR). In agreement with the above, pure zinc was found in extrusion to recrystallize dynamically above 0.6 Tm, while in creep this occurs above (0.7-0.8) 7^. In torsion, pure zinc shows evidence of dynamic recrystallization (DRX) between 0.6 Tm and 0.9 Tm (McQueen and Bourell, 1988). For a solid solution zinc alloy (Zn0.19%Cu-0.5%Cd-0.01%Mg), it was reported that, in extrusion, DRX takes place above 0.7 Tm, which illustrates the effect of alloying elements in solution on the recrystallization temperature (McQueen and Bourell, 1988). Zinc alloys also exhibit superplastic behavior with high strains and high strainrate sensitivities when the grain size is small. Boundary diffusion and sliding are considered to be the governing factors of this behavior (Alden, 1975). Superplastic behavior has been studied in numerous zinc alloys such as Zn-O.2%A1 or the
105
eutectoid alloy Zn-22%A1. A more comprehensive account of the superplastic behavior of metals is given by Mukherjee in Vol. 6 of this Series. 3.4.2 Alloys The main rolled alloys were introduced with the discovery of the effect of titanium additions by work at New Jersey Zinc Corporation in the 1940s (Anderson et al, 1944). Table 3-4 lists the composition of zinc alloys considered for rolling. Titanium induces the formation of an eutectic phase Zn-TiZn 1 5 . Titanium does not have a great effect on solid solution strengthening or the recrystallization temperature but the dispersion of the TiZn 15 provides strong barriers to grain growth and therefore produces grain boundary strengthening and increases the creep strength, the effect on the latter being dramatic (Rennhack and Conrad, 1966 a). Titanium has a strong structure refining effect in the cast condition by causing the nucleation of Zinc on Zn 2 TiO 4 particles forming by reaction of titanium with zinc and oxygen in the melt (Leone et al., 1975). Other elements forming eutectics with zinc, such as manganese, nickel or chromium, also give rise to a strengthening dispersion
Table 3-4. Wrought rolled zinc alloys (Goodwin and Ponikvar, 1989). Composition
Applications
Pure zinc Zn-08%Cu
Deep drawn hardware, expanded metal; Building construction materials, deep drawn hardware, coinage; Roofing, gutters, and down spouts, building construction materials, deep drawn hardware, address plates, solar collectors; Building construction materials, dry cell battery cans, deep drawn hardware, address plates, electrical components Shaped components such as typewriter castings, computer panels and covers; Photoengraving plates.
Zn-0.2%Cu-0.1%Ti Zn-0.07%Pb-0.05%Cd Zn-22%Al-0.5%Cu-0.01 %Mg Zn-0.1%Al-0.06%Mg
106
3 Zinc
of intermetallic-phase particles and therefore to improved creep strength (Pelzel, 1973). Copper in the contents shown (up to 1%), as well as cadmium, provides solid solution strengthening and increases the recrystallization temperature (Farge and Williams, 1965). Higher additions of copper increase strength further through additional solid solution strengthening and through precipitation of the s phase for copper contents above 2% (Pelzel and Paschen, 1962). In binary Zn-Mg, magnesium was shown to increase the recrystallization temperature (Farge and Williams, 1965) and to provide substantial creep strengthening for contents in the range of 0.001 to 0.0045% (Pelzel, 1970). In some ternary Zn-Cu-Mg alloys, magnesium in the range of 0.006 to 0.016% provides strengthening and stiffening (Mosbach, 1975). The solubility of magnesium in zinc being very low (less than 0.008% at 200 °C), it is the precipitated form of magnesium which acts on strength through a decrease of the subgrain size (Wegria and Habraken, 1970). The Zn-Al-Mg alloy for photoengraving plates (Table 3-4) is subject during thermomechanical treatment to aluminum segregations at the grain boundaries, which then control the behavior of the alloy during etching (Wegria et al., 1981). The composition of the alloys for battery cells (Table 3-4) is designed mainly for optimum electrochemical behavior of the battery: absence of passivation, uniform dissolution, etc. In terms of the fabricability, an easy inverse extrusion process is desired. The mechanics of the latter, especially as regards the role of lubrication, have been discussed by Magnee (1990). Research has been carried out in several organizations with the purpose of developing still stronger alloys for rolling or ex-
trusion (Crocq, 1976; Barnhurst, 1990 b). Alloys with substantial strength were identified as containing up to 30% aluminum or up to 7% copper. However, such alloys suffer from reduced ductility and their expansion has not been very large outside specific or temporary niches. 3.4.3 Casting and Rolling Processes
At present, typical processes used are as follows (Habraken, 1976; Coutsouradis, 1978): • continuous casting in belt- or roll-type machines followed by continuous or semicontinuous rolling, • casting in slabs and then rough rolling and finally continuous rolling. The main features of these processes relevant to the physical metallurgy of rolled alloys will be discussed in the next sections on the basis, unless otherwise indicated, of the review by Wegria (1985). 3.4.3.1 Casting
In belt-type continuous casting machines, the zinc alloy is cast in the form of thin slabs of, typically, 9 to 16 mm thickness at speeds of 5-15m/min and throughputs of up to 60 t/h. The as-cast microstructure of a Zn-Ti alloy consists of solid solution cells surrounded by the decomposed Zn/TiZn 15 eutectic (Rennhack and Conard, 1966 a). 3.4.3.2 Rolling
After the alloy has been cast in a belttype machine, rolling occurs depending on the design of the particular plant: the 9 18 mm thick slabs are rolled in 3 to 5 passes down to about 0.6 mm in one (Levasseur, 1976), two (Carpentier and Dreulle, 1972) or five roll stands (Anonymous, 1969; Wincierz, 1975).
3.4 Wrought Alloys
The rolling sequence is carried out in such a way that reheating does not occur during rolling. Similarly, there is no final intentional heat treatment except for the natural soaking of the coil at the final rolling temperature until cooling to room temperature. The rolling process causes several microstructural and texture effects. In the case of the Zn-Cu-Ti alloy, the breakdown of the eutetic results in the dispersion of the TiZn 15 compound generally in the form of small particles (0.1 ^im) aligned in filaments or fibers and originating from the initial as-cast microstructure (Rennhack and Conard, 1966 a). The dispersed TiZn 15 phase is essential since it pins the grain boundaries and thus prevents grain boundary migration and, therefore, excessive grain growth (Rennhack and Conard, 1966 b). Typical microstructures of a rolled Z n Cu-Ti alloy are reproduced in Fig. 3-8 (European Zinc Producers, 1981), showing recrystallized grains and the presence of a sub-structure. The rolling process causes microstructural modifications. High temperatures for the final rolling pass favor increased creep strength while lower temperatures for the final passes lead to better ductility (Wincierz, 1975). The rolling process also determines the formation of texture. The latter is responsible for the isotropy or lack of isotropy of the mechanical properties - especially ductility (Burggraf and Wincierz, 1981) - and physical properties. In particular, the coefficient of thermal expansion depends on texture. Specific rolling conditions, i.e. heavy deformation at elevated temperatures, result in the formation of a c-axis texture in which the basal plane of the hexagonal crystal is in the plane of the sheet (Cosse etal., 1980).
107
(a)
(b)
Figure 3-8. Microstructure of a rolled Zn-Cu-Ti alloy: (a) SEM, (b) TEM.
3.4.4 Properties
Table 3-5 presents typical properties for the Zn-0.2%Cu-0.1%Ti sheet alloy. Owing to the heterogenous microstructure and to texture effects, the mechanical properties are not isotropic. The creep deformation of Zn-Cu-Ti alloys was shown to be represented by various creep correlating parameters (Cosse
108
3 Zinc
Table 3-5. Typical RT mechanical and physical properties of the Zn-Cu-Ti sheet alloy. Property
LongiTranstudinal verse direction direction
Tensile strength in MPa 155-170 190-210 Yield strength in MPa 120-140 150-170 Elongation in % 40- 60 20- 40 Young's modulus in MPa 61000 86000 53-58 Hardness in HV3 Creep rate at RT under 70 MPa in %/h 10" Specific weight in g/cm3 7.14 Coefficient of linear expansion inK"1 20x10"
et al., 1980), allowing extrapolation up to a factor of about 30. The deformation behavior in drawing of zinc alloy sheet (Zn-2.5%Cu-0.15%Ti and Zn-0.2%Cu-0.1%Ti) was shown to be governed by different properties compared to other materials such as steel (Cosse et al., 1978). In the case of zinc sheet materials, it was shown that since n, the strain sensitivity, is small, the Erichsen value is related to the parameter n + 5 m, m being the strain rate sensitivity. Bending is an important property for sheet because roofing components are fabricated essentially by bending, often at small angles. Bending depends on various microstructural features and on texture. The latter allows a favorable orientation for pyramidal glide of the 2d order (Wegria, 1985; Galledou, 1992).
3.5 Recycling In the production of zinc in the western world, the use of scrap and residues accounted in 1990 for 22.9% of the total consumption of 6.25 Mt (Metallstatistik, 1991). This percentage is lower than for metals
such as lead (52%), copper (40%) or even aluminum (27%). A probable reason is that one of the main uses of zinc is in association with steel as a coating with long service periods. Particular mention must be made here of the problem area of scrap of galvanized steel combined with steel scrap, which is mainly a result of the extensive use of coated steel in cars. Although there are many processes under study, we simply mention the use of zinciferrous steel scrap in electric arc furnaces (EAFs) (Hofiken et al., 1992), the subsequent treatment of the dusts in Waelztype furnaces and the final treatment of the concentrates of zinc-lead in primary zinc smelters using the imperial smelting process. Reviews on the subject are available (Bounds, 1990).
3.6 Concluding Remarks Some forms of products which are well established for many metals, such as powder metallurgy, extrusion for sections or tubes, closed die forging etc., are not important in zinc. This situation may be explained by the existence of alternative possibilities for zinc, e.g., die-casting, which makes some of the advantages of processes such as powder metallurgy or closed die precision forging less evident. Most of the processes developed during the last two decades for forming and reinforcing of metals are applicable to zinc and zinc alloys. Such processes include squeeze casting, stir casting, particulate or fiber composites by casting or powder metallurgy routes, spray deposition, etc. Results on such processes applied to zinc alloys were reviewed by Porter (1991). The zinc alloys discussed in previous sections are well established and it appears that, aside from exploratory work in new
3.8 References
directions, the general trend for research is not in the development of completely new alloys but rather in better and consistent uses of the existing ones. In continuous hot dip galvanizing, the alloys already available (zinc, galvannealed, Galfan, Galvalume) provide an ample field for a better understanding of their behavior, improved processing, modeling and establishment of data bases. New processes, e.g., vapor deposition, when established, may well result in entirely new systems for the protection of steel. The same trend is also valid for galvanizing after manufacturing, although the discontinuous nature of the process requires special approaches. In terms of alloys, a process for systems other than zinc, e.g., Zn-5%A1, has not as yet been developed for widespread use in applications, such as transport, for which a reduced weight of coating is desirable while maintaining a high corrosion resistance. In the area of the protection of steel, the intermaterial competition is low, zinc coatings providing a most economical solution. The casting alloys provide a wide range of compositions, the near eutectic Zn-4%A1 alloy being predominant in diecasting. Further research is being carried out in thin wall pressure die-casting in order to better resist the fierce competition especially from plastics and aluminum. Rolled alloys appear to be well established and their use in architecture has led to great advances. Areas of current research include deformation behavior as a function of textures and microstructures. Competition from plastics also here is severe, while competition from other metals (e.g., copper, stainless steel) is only marginal. Another direction of expanded or more efficient use of zinc alloys may be in the form of combinations of different materi-
109
als, for example, as in the case of components consisting of steel, zinc (coatings) and polymers.
3.7 Acknowledgement The authors acknowledge the Zinc Committee of CRM (Asturienne des Mines, Metallurgie Hoboken Overpelt, Vieille Montagne, now consolidated in Union Miniere) for support in a number of research projects over many years.
3.8 References Alden, T. H. (1975), Treatise on Materials Science and Technology, Vol. 6: Plastic Deformation of Materials. San Francisco: Academic Press, p. 225. Anderson, E. A., Boyle, E. J., Ramsey, P. W. (1944), Trans. AIME 156, 278. Anonymous (1969), Metall 23, 857. ASM (1990), ASM Handbook, Vol. 2: Zinc and Zinc Alloys, 10th ed. Metals Park, OH: ASM. ASM (1992), ASM Handbook, Vol. 3: Alloy Phase Diagrams. Metals Park, OH: ASM. Barnhurst, R. I (1990 a), / Mater. Eng. 12 (4), 279. Barnhurst, R. J. (1990 b), Metals Handbook, Vol. 2: Zinc and Zinc Alloys, 10th ed. Metals Park, OHr ASM, p. 526. Barnhurst, R. J. (1992), Die Casting Engineer, Jan.Feb., 32. Barnhurst, R. I (1993), in: 14th Int. Pressure Die Casting Conf, May 4-5, 1993, Birmingham, England. London: Zinc Development Association. Bounds, C. O. (1990), "Recycling of Zinc from EAF Dust", in: ASM Handbook, Vol. 2, 10th ed. Metals Park, OH: ASM, p. 1224. Burggraf, X, Wincierz, P. (1981), Z. Metallkd. 72, 287. Burkhardt, A. (1940), Technologie der Zinklegierungen. Berlin: Springer. Carpentier, A., Dreulle, N. (1972), Metall Rep. CRM 33, 73. Christman, T. K., Payer, I, Kurr, G. W, Doe, J. B., Carr, D. S., Ponikvar, A. L. (1986), Zinc: Its Corrosion Resistance. Research Triangle Park, USA: ILZRO. Cosse, P., D'Haeyer, R., Coutsouradis, D., Habraken, L. (1978), Mem. Sci. Rev. Metall., Aug.) Sept., 517. Cosse, P., Coutsouradis, D., Habraken, L., Piccinin, A. (1980), Z. Metallkd. 71, 138. Coutsouradis, D. (1978), ATB Metall. 18, 165.
110
3 Zinc
Coutsouradis, D., Ayoub, C. (1989), Metall Ital. 81, 605. Coutsouradis, D., Goodwin, F.E., Pelerin, X, Skenazi, A. F. (1984), Stahl, Eisen 104, 1073. Crocq, G. (1976), ATB Metall XVI, 201. Davies, R. W, Williams, P. (1975), Met. Sci. 9, 149. Dreulle, N. (1979), in: Intergalva 1979, 12th Int. Galvanizing Conf., Paris, France. London: Zinc Development Association. Dupouy, J.M., Poirier, J. P. (1968), in: Deformation plastique des metaux et alliages. Paris: Masson, PP.
l-n.
European Zinc Producers (1981), Metallographic Atlas of Zinc and Zinc Alloys. Liege: CRM. Farge, J. C. T., Williams, W M. (1965), Can. Metall. Q. 5, 265. Galledou, Y. (1992), Contribution a Vetude de la pliabilite des alliages Zn-Cu-Ti a basse temperature. D.Sc. Thesis, University of Metz. Galvez-Soza, D., Greday, T., Coutsouradis, D. (1984), ATB Metall. 24, 253. Gervais, E., Barnhurst, R. X, Loong, C. A. (1985), /. Met. Nov., 43. Goodwin, F. E., Ponikvar, A. L. (1989), Engineering Properties of Zinc Alloys, 3rd ed. Research Triangle Park, USA: ILZRO. Goodwin, F. E., Walmag, G., Murphy, S., Wegria, X (1993), in: 14th Int. Pressure Die Casting Conf, May 4-5, 1993, Birmingham, England. London: Zinc Development Association. Habraken, L. (1976), ATB Metall. 16, 176. Hisamatsu, Y. (1989), in: Galvatech 89, Int. Conf on Zinc and Zinc Alloy Coated Steel Sheet, Tokyo: The Iron and Steel Institute of Japan, p. 3. Hoffken, E., Wolf, X, Kuhn, P. (1992), Thyssen Tech. Ber. No. 1, 119. Horstmann, D. (1978), The Reactions between Iron and Liquid Zinc. London: Zinc Development Association. Humphreys, X F. (1981), in: Proc. 2nd Int. Riso Symp.: Hansen, N. (Ed.). Ris0, Denmark: Riso National Laboratories, p. 35. Jacquerie, X M. (1966), Metall. Rep. CRM 9, 51. Jaoul, B. (1965), Etude de la plasticite et application aux metaux. Paris: Dunod. Kingsbury, G. R. (1992), Metals Handbook, Vol. 18: Friction and Wear of Sliding Bearing Materials, 10th ed. Metals Park, OH: ASM, p. 741. Krepski, R. P. (1980), The Influence of Bath Alloy Additions in Hot-Dip Galvanizing - A Review. Monaco, PA: St Joe Minerals Corporation. Kriner, S. A., Niederstein, K. (1991), in: Intergalva 91, 3rd Int. Zinc Coated Sheet Conf, Barcelona, June 1991. London: Zinc Development Association, Paper S1A/1. Lamberigts, M., Walmag, G., Coutsouradis, D., Delneuville, P., Meeus, M. (1985), 89th Casting Congress AFS, April 29-May 3, 1985. Pittsburgh, PA: American Foundrymen's Society, Paper 85-84, 569.
Lamberigts, M., Leroy, V., Beguin, M., Dubois, M., van der Heiden, A., Verhoeven, G. (1992), in: Galvatech 1992, 2nd Int. Conf on Zinc and Zinc Alloy Coated Steel Sheet. Dusseldorf: Verlag Stahleisen, p. 199. Lecomte-Beckers, X, Terzief, L., Rashev, P., Walmag, G. (1989), /. Microsc. Spectrosc. Electron. 14, 85. Leone, G. L., Niessen, P., Keer, H. W (1975), Metall. Trans. B6, 503. Leroy, V., Pelerin, X, Emond, C , Habraken, L. (1976), "Galvanizing of Silicon-Containing Steels", in: Proc. 2nd Int. ILZRO Galvanizing Seminar, Saint-Louis, MO. Research Triangle Park, NC: ILZRO, Levasseur, X (1976), ATB Metall. 16, 180. Magnee, A. (1990), Eur. J. Mech. Eng. 35, 155. Masson, N. (1991), ATB Metall. 21, (3-4), 44. McQueen, H. X, Bourell, D. L. (1988), /. Mater. Shaping Technol. 5, 163. Metallstatistik 1980-1990 (1991) 78th ed. Frankfurt am Main, FRG: Metallgesellschaft. Meuthen, B. (1989), in: Galvatech 89, Proc. Int. Conf on Zinc and Zinc Alloy Coated Steel Sheet. Tokyo: The Iron and Steel Institute of Japan, p. 91. Michel, X P. (1979), Influence d'unepredeformation en glissement pyramidal sur le glissement basal du zinc. Ph.D. Thesis, University of Nancy. Mosbach, W. X (1975), Met. Prog., Feb., 60. Niessen, P., Wineguard, W. C. (1966), / Inst. Met. 94, 31. Ogier, G. (1993), in: 14th Int. Pressure Die-Casting Conf, Birmingham, UK, May 1993. London: Zinc Development Association. Partridge, P. G. (1967), Int Metall. Rev. 12, 169. Pelerin, X (1983), Contribution a Vetude de la reaction Fe-Zn. Ph.D. Thesis, University of Lille. Pelerin, X, Coutsouradis, D., Foct, X (1985), Mem. Sci. Rev. Metall. April, 191. Pelzel, E. (1970), Metall 24, 961. Pelzel, E. (1973), Metall 27, 106. Pelzel, E., Paschen, P. (1962), Metall 16, 750. Porter, F. (1991), Zinc Handbook: Properties, Processing and Use in Design. New York: Marcel Dekker. Radtke, S. F. (1972), Metall 28, 891. Renaux, B., Silvestre, A. (1988), in: Intergalva 88, Proc. 2nd Int. Zinc. Coated Sheet Conf, Rome, June 1988. London: Zinc Development Association, Paper SB3/1. Rennhack, E. H., Conard, G. P. (1966a), Trans. Metall. Soc. AIME 236, 694. Rennhack, E. H., Conard, G. P. (1966b), Trans. Metall. Soc. AIME 236, 1441. Sandelin, R. W. (1940), Wire, Wire Prod. 15, 655. Schwarz, W, Schumacher, B., Wolfhard, D. (1992), in: Galvatech 1992, Proc. 2nd Int. Conf. on Zinc and Zinc Alloy Coated Steel Sheet. Dusseldorf: Verlag Stahleisen, p. 417. Skenazi, A. F , Pelerin, X, Coutsouradis, D., Magnus, B., Meeus, M. (1983), Metall 37, 898.
3.8 References
Sokolowski, R. (1988), in: Intergaha 88, Proc. 15th Int. Galvanizing Conf., Rome, Italy. London: Zinc Development Association, Paper GEL Spittle, J. A., Brown, S. G. R. (1993), in: 14th Int. Pressure Die Casting Conf., Birmingham, UK, May 1993. London: Zinc Development Association. Wajda, E. S. (1954), Acta Metall. 2, 184. Walmag, G., Lamberigts, M., Coutsouradis, D. (1987), Unpublished data, Belgian Patent, No. 087.00555 (13 May 1987). Walmag, G., Murphy, S., Skenazi, A., Goodwin, F. (1991), in: Proc. of the 16th NADCA Die Casting Congress and Exposition, September 30-October 3, Detroit, ML River Grove, IL: NADCA. Wegria, J. (1985), Etude de la plasticite'des alliages de zinc—cuivre—titane. Application a Vamelioration de leur aptitute au pliage. Ph.D. Thesis, University of Lille. Wegria, X, Habraken, L. (1970), Metall Rep. CRM 22, 49. Wegria, I, Piccinin, A., Racek, R. (1981), Metall. Rep. CRM 58, 31. Wincierz, P. (1975), Z. Metallkd. 66, 235.
111
General Reading Barnhurst, R. J. (1990), Metals Handbook, Vol. 2: Zinc and Zinc Alloys, 10th ed. Metals Park, OH: ASM, p. 526. Burkhardt, A. (1940), Technologie der Zinklegierungen. Berlin: Springer. Horvick, E. W. (1979), Metals Handbook, Vol. 2: Selection and Application of Zinc and Zinc Alloys, 9th ed. Metals Park, OH: ASM, p. 629. Kingsbury, G. R. (1992), Metals Handbook, Vol. 18: Friction and Wear of Sliding Bearing Materials, 10th ed. Metals Park, OH: ASM, p. 741. Porter, F. (1991), Zinc Handbook: Properties, Processing and Use in Design. New York: Marcel Dekker. Smithells, C. J. (1976), Metals Reference Book. London: Butterworths.
4 Magnesium-Based Alloys Giinter Neite Metallgesellschaft AG, Frankfurt, Germany Kohei Kubota Mitsui Kinzoku, Saitama, Japan Kenji Higashi University of Osaka Prefecture, Osaka, Japan Franz Hehmann ONERA, Chatillon, France
List of 4.1 4.2 4.2.1 4.2.2 4.2.2.1 4.2.2.2 4.2.2.3 4.2.3 4.2.4 4.2.5 4.3 4.3.1 4.3.1.1 4.3.1.2 4.3.2 4.3.2.1 4.3.2.2 4.3.2.3 4.3.2.4 4.4 4.4.1 4.4.2 4.4.2.1 4.4.2.2
Symbols and Abbreviations Magnesium at the Turn of the Millennium Characteristics of Magnesium Magnesium Metal Structural Criteria Related to Mechanical Properties Mechanical Particulars: Magnesium Versus Aluminum Rate-Controlling Dislocation Reaction Effect of Grain Size and Thermal Activation on Deformation Modes Corrosion Behavior of Magnesium Further Properties of Engineering Relevance Summary of the Advantages and Limitations Magnesium Alloys and Their Applications Equilibrium Alloying Magnesium Casting Alloys Magnesium Alloys for Elevated Temperature Application Wrought Magnesium Alloys Processing Mechanical Properties at Room Temperature Mechanical Properties at Elevated Temperatures Magnesium-Lithium Alloys Rapid Solidification Processing of Wrought Magnesium Alloys Synopsis Brief Summary Atomization Versus Chill-Block Methods State of the Art and Alloy Availability
Materials Science and Technology Copyright © WILEY-VCH Verlag GmbH & Co KGaA. Allrightsreserved.
115 117 120 120 121 121 124 126 128 130 131 132 132 134 143 150 151 152 155 157 161 161 163 163 165
114
4.4.3 4.4.3.1 4.4.3.2 4.4.4 4.4.4.1 4.4.4.2 4.4.4.3 4.4.5 4.4.6 4.4.6.1 4.4.6.2 4.4.7 4.4.7.1 4.4.7.2 4.4.8 4.4.8.1 4.4.8.2 4.4.9 4.4.9.1 4.4.9.2 4.4.9.3 4.5 4.6 4.7 4.8 4.9
4 Magnesium-Based Alloys
Chill-Block Quenched Magnesium-Based Engineering Alloys: Extrusions . High Strength and High Corrosion Resistance at Ambient Temperatures .. High Temperature Properties Near-Net-Shape Forming of Chill-Block Quenched Magnesium Alloys . . . Forging of Monolithic EA55RS, the Resultant SiCp Composites, and RS AZ91 Sheet Rolling of EA55B-RS Superplasticity Gas-Atomized and Spray-Formed Magnesium-Based Engineering Alloys . Rapid Solidification Processed (RSP) Magnesium-Lithium Alloys State of Development Magnesium-Lithium-Hydrogen Alloys Amorphous Magnesium Alloys Amorphization by Liquid Quenching Solid State Amorphization Hypersaturated Solid Solutions Passivation of oc-Magnesium Novel Constitutional Thresholds for Enhanced Response to Ageing Development Principles for Rapid Solidification Processed (RSP) Wrought Magnesium Alloys Damage Tolerance by "Clean" Grain Refinement Effect of Grain Refinement on Hall-Petch Strengthening Multiphase Dispersions Magnesium Metal Matrix Composites Recycling of Magnesium Alloys Conclusions and Outlook Acknowledgements References
169 169 173 175 175 176 177 178 181 181 183 185 186 187 188 190 191 193 193 194 196 198 202 203 206 206
List of Symbols and Abbreviations
List of Symbols and Abbreviations a,c b d E h K kH
K L m n T
h TEx Tm ATN Tx
K ^NHE
lattice constants Burgers vector droplet size; fragmented size; critical thickness; grain size Young's modulus elongation-to-fracture heat transfer coefficient shear modulus coupling factor for UTS, UTS = (fcH x microhardness) Hall-Petch coefficient impact strength length strain-rate-sensitivity exponent order of magnitude of the production scale temperature intrinsic fault vector extrusion temperature melting temperature difference in undercooling crystallization temperature externally controlled growth rate internally controlled growth rate standard potential against normal hydrogen electrode recalesced solidification velocity
a y, yb yp £ e 6 k Q
coefficient of thermal expansion stacking-fault energy prismatic stacking-fault energy strain strain rate wetting angle angle which the slip direction makes with the tensile axis density; electrical resistivity
initial stress rotating bending fatigue strength yield stress critical resolved shear stress angle between the normal to the slip plane and the tensile axis diameter
AE b.c.c. BHN crss CTE
alkaline earth body centered cubic Brinell hardness number critical resolved shear stress coefficient of thermal expansion
115
116
CVD CYS DC DSC EBED EBEP f.c.c. GA GBS GP Gr h.c.p. HP I/M LDC MM MMC mpy MS NHE PFC PM PVD RE ROM RS RSP SCE SF TEM TM TR TSSE TYS UTS VHN
4 Magnesium-Based Alloys
chemical vapor deposition compressive yield strength direct current differential scanning calorimetry electron beam evaporated deposit electron beam evaporated powder face centered cubic gas atomization/atomized grain boundary sliding Guinier-Preston graphite fiber reinforced hexagonal close-packed high purity ingot metallurgy liquid dynamic compaction mischmetall metal matrix composite mils per year melt spinning/spun normal hydrogen electrode planar flow casting powder metallurgy physical vapor deposition rare earth rule of mixtures rapid solidification rapid solidification processing/processed saturated calomel electrode stacking fault; spray formed transmission electron microscopy/microscope transition metal twin-roller terminal solid solubility extension tensile yield strength ultimate tensile strength Vickers hardness number
117
4.1 Magnesium at the Turn of the Millennium
4.1 Magnesium at the Turn of the Millennium Sustainable development of light structural materials did not occur until the twentieth century (Fig. 4-1). Unprecedented growth rates of around 0.4 orders of magnitude per decade for aluminum and plastics illustrate their versatility for mankind's welfare. The development of heavy materials, following their introduction in the 19th century or earlier, was much less rapid, for good reasons (see Thorbeck and Hehmann, 1996; Fig. 4-2). Progress in extraction, reduction, and refining technologies has since assured sustainability in virtually all sectors of the production of structural materials, without particular need to address fuel efficiency, for example. The increase in the average growth rate with decreasing density and toxicity, for example, indicates that mobility, lightness, and also social urbanization and individual health including fitness for use in cities - are becoming the major stimuli in the selection of new, alternative materials at the end of the second millennium. While aluminum and plastics show typical Johnson-MehlAvrami growth kinetics of production, magnesium remains quite different in this respect. While it is fashionable to admire the potential role of new materials, it is even more intriguing to see what is missed out as time goes by. Ceaseless industrial criticism is at least as counterproductive as widespread optimism in forecasting the role of new materials. Key concepts derived from sustainable development principles should thus increase the appeal of magnesium (Thorbeck and Hehmann, 1996). The use of magnesium has become inevitable in the modern world, or: the crisis itself triggers the boom. What are the conditions for the boom? What are the possibilities offered to us? Magnesium has high
109 Light Structural Materials in the 20th Century 108
1900
1920
1940
1960
1980
2000
Year Figure 4-1. Production of aluminum, magnesium, and plastics during the 20th century. Data for 19901992 are highlighted by a circle [after Metallstatistik (1990); supplementary data from Metallgesellschaft, Frankfurt, for aluminum and magnesium, and from Verband der Kunststofferzeugenden Industrie e.V. (VKE, 1993), Frankfurt, for the plastics]. n, order of magnitude of the production scale. 1} After Metallstatistik (1990), including data from the U.S.A., Japan, and Western Europe without Switzerland, the Netherlands, Spain, and Yugoslavia. Note that the former U.S.S.R. (C.I.S.) covers about 40% of the corresponding markets; 2) after International Magnesium Association (1993), without data for former U.S.S.R. (C.I.S.) and P.R. China.
supranational availability (earth's crust: 2 mass% or rank 6; sea water: 1.1 kg/m 3 or rank 3 among the dissolved minerals) and unusually varied industrial uses, with key positions as a constituent in aluminum, or for steel, chemical, and pharmaceutical products, not forgetting the continuous intake by mammals, including that of the human species, which is now approaching 0.4 g/day or so. The potential (driving forces) for substantial growth using mag-
118
4 Magnesium-Based Alloys 8n 10 years
IB 2-2 ®
Mg (1932 - 1942)
>
3
extrapolated 'metallic" sustainability potential
Plastics 0.4
-
0.2
_-
© ©
0.0
I
castMg^^-^
Zn Fe Cu ^ — ^ •—-
wrought Mg I
l
i
Pb
^ — • 1
i
10
12
Density p ( g/cm3) Figure 4-2. Mean growth exponent for production per decade, 5«(10), of light and heavy materials during the 20th century (solid circles) versus density, indicating that the lighter the material, the higher the average growth rate. 8n(10) = 10[log(P2)-log(Pl)]/[Yr(2)- Yr(l)], where P2 is the production of the metal concerned in the last year of available statistics, PI is the production of the metal in the first year of available statistics, Yr(2) is the last year of available statistics and Yr(l) is the first year of available statistics. For comparison, corresponding growth rates for wrought magnesium alloys, cast magnesium alloys, and magnesium production from 1932-1942 are included. Except for the latter, the time integral is from the first available year with an overall production of more than 1000 metric tons until the last recorded year (i.e. 1990 or 1992). The comparison shows how multi-industrial use (by independent branches of industry) of primary magnesium might obscure the outlook magnesium alloys as well as the potential of magnesium for ultra-rapid growth, as was evident in 1932 to 1942 when the overall growth was the prime result of the use of corresponding alloys. Sources as for Fig. 4-1 and Metallstatistik (1990) for Zn, Fe, Cu, and Pb.
nesium is not negligible if considered seriously. Once the logarithmic scale has been accepted as the universal scale for growth, the other side of the law has to be acknowledged also, that is, the lack of dependence of the growth rate on the absolute scale, including the possibility of mutation such as abrupt leaps caused by man's impact on political, economic, and natural frontiers. Nazi Germany's war needs triggered the growth of magnesium in the 1930s. The growth rate of magnesium production at that time has not been surpassed by any structural material since then, but the impact the war had on sustainability was practically nil (see Fig. 4-2 and compare the aluminum curve after World War II in
Fig. 4-1). It is, however, also evident that it was not peace which ended this development, but the absence of a self-healing, naturally passivating surface film. This situation is reminiscent of that faced by the steel industry some hundred years ago. During war, the most important question is "Who will win the war?" During peace, the focus shifts to concepts for maintaining peace. However, even during peace a true equilibrium is never reached. Bombs have been replaced by pollution, and the use of magnesium offers a chance to reduce this pollution significantly. Any global topic has a political dimension and any political sustainability builds upon economic sense. The potential for replacing aluminum castings by magnesium
4.1 Magnesium at the Turn of the Millennium
castings, for example, corresponds to 1.5 orders of magnitude, but the corresponding and more neglected potential for magnesium wrought products is almost 4 orders of magnitude (Fig. 4-1). The wrought magnesium market is practically untouched. If the potentially replaceable plastic-based products are added, the relative growth potential appears to be approximately as large as the universe (see Emley, 1966). The small extent of today's magnesium alloy market obscures the fact that magnesium is the structural material of the future for both airborne and terrestrial transport. An early indication of this fact is the on-going growth of the magnesium alloy market during the early 1990s despite worldwide recession, thus making magnesium one of the few structural materials with a growing market at that time (Fig. 4-1, circles at bottom). The actual growth, however, originates only from the introduction of high-purity castings, which dates back to the early 1980s. The introduction of magnesium has been plagued by counterproductive attitudes. Emphasis on primary magnesium raw material production cannot circumvent the need to arouse greater interest
119
from the alloy user side, while emphasis on automobile applications cannot circumvent the need to regain the aerospace industry as the workhorse for sustainable development, including the spin-offs that are bound to come. Nevertheless, nowhere in the world of materials are possibilities and reality turned upside down so much as in connection with magnesium. About 80% of the magnesium manufactured is used as an additive for aluminum alloying or for extractive metallurgy, for example, while 13% is used for casting alloys and 3% remains for the more promising wrought products (Fig. 4-3). When it comes to the global perspective, nowhere else are science and research on the one hand and production on the other drifting apart so strongly. This chapter provides information on novel methods of production. We hope to supply readers with a foundation that will enable them to identify their own concrete area(s) for growth. We do not discuss alloy development approaches, since an approach alone is not sufficient to ensure success. The obscured but rich diversity, selfdynamism, and flexibility of magnesium alloy development at the onset of the next
Total demand
Aluminium alloying
257 300 t
Other Wrought products Metal reduction Nodular iron Die casting
Figure 4-3. Primary magnesium conChem./electrochem. use Desulfurization
sumption data for 1992 (Garber, 1993).
120
4 Magnesium-Based Alloys
millennium will assure sustainability to counteract the current recession and also the realization of unexpected properties and productivities, although it is as yet impossible to comprehend the significance and implications for science, industry, and society. Magnesium will stand for the special qualities of the third millennium. There will be less emphasis on the availability of materials and more emphasis on sophisticated use.
200
> 100
o O
0 400
W0
300
200
4.2 Characteristics of Magnesium
100
NU-
••
4.2.1 Magnesium Metal Mg
Magnesium metal can be obtained from magnesite, carnallite, or dolomite ore, for example, from salt deposits, from natural underground and artificial brines, as well as from seawater. The prime reduction principle is the electrolysis of molten magnesium chloride (MgCl2), derived chemically from dolomite by calcination, precipitation, and chlorination, or directly from brines, followed by the thermal reduction of magnesium oxide (MgO) (Kirk-Othmer, 1981). The electrolysis of MgCl2 should require an electric energy of 25 MJ/kg, but current methods need 40-80 MJ/kg due to recombination effects and limited current efficiency in the electrolyte cell. All major producers have or intend to raise the efficiency of their process to achieve a value of 40 MJ/kg or less. For comparison, 47 MJ/kg is needed for producing aluminum metal from A12O3 by electrolysis. The total energy required from ore extraction to primary metal is shown in Fig. 4-4 for different materials (Muller, 1991; Kindler and Nikles, 1980). For primary magnesium a total energy of about 125 MJ/kg is necessary. On a volume basis, however, magnesium production requires less en-
Al
Zn
ABS PMMA Polyamide
Figure 4-4. Total energy required on the basis of (top) weight and (bottom) volume for producing different primary materials [after Muller (1991) for the metals and Kindler and Nikles (1980) for the plastics].
ergy than aluminum production. It should be noted that the potential for further reduction in the energy consumption is greater for the production of magnesium than for aluminum. Due to magnesium's low heat content less energy is required to achieve the electrolysis process temperature, more than compensating for the slightly higher binding energy. The demand for raw magnesium metal differs significantly from that of other raw structural metals in that about 80% is used for competing materials, such as for alloying aluminum, desulfurizing steel, and the reduction of titanium (Fig. 4-3). The major areas of growth, however, are not within these applications, but in structural applications of magnesium alloys (see also Sees. 4.1 and 4.7). This has stimulated recent efforts to enhance established alloy families such as Mg-Al-Zn-based alloys
4.2 Characteristics of Magnesium
using conventional and more advanced technologies, as well as to develop entirely new alloys. Our intention is not only to review the reported engineering properties of new magnesium alloys, but also to refer to the incomplete and still immature developments. These developments are directed towards the more effective exploration of the basic properties of magnesium and the progress in processing technologies made over the last few decades. The average magnesium enthusiast will find a summary of its basic properties in Table 4-1. As a group IIA alkaline earth metal situated between beryllium and calcium, magnesium is divalent and shiny silver in color, has a density of 1736 kg/m 3 at 25 °C as derived from X-ray analysis, and a hexagonal close-packed crystal structure of which the ratio of the crystal axes c and a is 1.6236 (i.e., not very far from the ideal value for dense atomic packing, 1.633). This crystal structure is one of the factors which lead to the inferior mechanical properties of common commercial magnesium alloys (whose densities do not exceed 1900 kg/m3) compared to b.c.c. iron or fee. aluminum, for example, but it should also be noted that the effect of higher order next-neighbors upon the mechanical properties of h.c.p. magnesium is far from understood (see Sees. 4.2.2 and 4.4.9). Evidently, the corrosion behavior is a key factor in the use of magnesium for structural applications (see Sec. 4.1). This corrosion behavior is discussed in Sec. 4.2.3 followed by a short introduction covering a number of basic advantages of magnesium which are easily overlooked (Sec. 4.2.4).
121
4.2.2 Structural Criteria Related to Mechanical Properties 4.2.2.1 Mechanical Particulars: Magnesium Versus Aluminum
Magnesium has attracted much attention for studying the nature of ideal metallic or free electron bonding (Raynor, 1959). The free 3 s2 valence electron structure of pure magnesium excludes it from any covalent bonding phenomena, resulting in the lowest average valence electron binding energies and the weakest interatomic cohesion of any structural metal (Kittel, 1976). The additional covalent 3px bond of pure aluminum results in larger moduli of elasticity E=7l GPa and shear K=26 GPa, and a larger specific modulus of elasticity E/Q = 27 GPa cm3/g than for magnesium (£=45 GPa, # = 1 7 GPa, £/£ = 25GPacm 3 /g; Table 4-1; Cottrell, 1975). The operative slip distance in magnesium is as for aluminum, but the perpendicular distance between two adjacent slip planes is about four times larger. Both the theoretical and the operative PeierlsNabarro stress on the close-packed {0001} planes of pure magnesium are thus orders of magnitude lower than on the {111} planes of pure aluminum. Perfect lattice dislocations in pure magnesium and conventional magnesium alloys would therefore have a relatively wide strain field and a much higher mobility than in pure aluminum and aluminum alloys. Despite this higher mobility, however, pure (99.95%) polycrystalline magnesium is harder [37 Brinell hardness number (BHN) at ambient temperature; Bollenrath, 1939] than pure (99.99%) polycrystalline aluminum (15 BHN at ambient temperature; Aluminium Taschenbuch, 1984). The HallPetch coefficient ky = 280 MPa / for magnesium is four times larger than that of
122
4 Magnesium-Based Alloys
Table 4-1. Physical and engineering properties of pure magnesium. Temperature in °C
Property
Atomic number Relative atomic mass Natural isotopes
Melting point Boiling point First ionization energy Structure a c c/a Density Electrical resistivity (polycrystalline) Elastic moduli Cll
25
25 20 600 25
c33 Q4
c12 c13
Young's modulus of polycrystalline Mg Poisson's ratio of polycrystalline Mg Coeff. of thermal expansion parallel to a parallel to c polycrystalline Linear contraction Volume contraction liquid-solid Heat capacity Cp Entropy S Enthalpy H-H25oC Thermal conductivity Thermal diffusivity Electrochemical potential (Normal hydrogen electrode) Rel. machining power Mg alloy :A1 alloy Mg alloy:cast iron Mg alloy :Ni alloy
25 25 27 527 27 527 27 527 650-20 °C 650 °C 27 527 27 527 527 27 527 27
Value
12 24.3050 79% HMg 10% l52Mg 1 1 % l62Mg (650.0 ± 0.5) °C 1090°C 7.646 eV hexagonal (hP2) 0.32094 nm 0.52107 nm 1.6236 1736 kg/m3 4.46xlO" 8 Qm 17.0xl0"8Qm 59.3 GPa 61.5 GPa 16.4 GPa 25.7 GPa 21.4 GPa 45 GPa 0.35 24.7xlO" 6 /K 29.8xlO" 6 /K 25.7xlO~ 6 /K 3O.5xlO~6/K 25.0xl0" 6 /K 30.0xl0~7K 1.9% 4.2% 24.86 J/mol K 31.05 J/molK 32.52 J/mol K 59.72 J/mol K 14057 J/mol 156W/mK 146 W/m K 0.874 cm2/s -2.37 V 1:1.8 1:3.5 1:10
Reference
Fluck and Heumann (1986) Fluck and Heumann (1986)
Massalski et al. (1990) Massalski et al. (1990) Fluck and Heumann (1986) Massalski et al. (1990)
from structure data Kirk-Othmer (1981) Landolt-Bornstein (1979)
Kirk-Othmer (1981) Kirk-Othmer (1981) Touloukian et al. (1978)
Kirk-Othmer (1981) Stull and Sinke (1956) Stull and Sinke (1956) Stull and Sinke (1956) Touloukian et al. (1978) Touloukian et al. (1978) Froats et al. (1987) Kirk-Othmer (1981)
4.2 Characteristics of Magnesium
123
1. One Basal Slip Plane (0001) a2 = aa=
(0002)
Type-II Planes
Type-I Planes
2. Two Prismatic Slip Planes
Mg-Li Mg-In
b=l/3a<1120:
7 (1010)
(1120)
3. Four Pyramidal Slip Planes Two First Order
(10T1)
high-Tslip
(1121)
Two Second Order
twinning
(1012)
b = l/3a+c<1123>
Figure 4-5. Slip planes and slip directions of h.c.p. magnesium in order of thermal activatability above ambient temperature {italics: conditions for operative slip model) (Emley, 1966).
(1122)
pure aluminum (68MPa v /m; Courtney, 1990). Below 498 K, plastic deformation of polycrystalline h.c.p. Mg is limited to basal {0001}<1120> slip and to pyramidal {1012}<10ll> twinning (Fig. 4-5; Emley, 1966). The magnesium crystal thus has only three geometrical and two independent slip systems, while the aluminum crystal has twelve {111}<1IO> geometrical and five independent slip systems for general shape change (von Mises, 1928). The maximum (polycrystal: mean) value of (cos(/> cos A), where (f) is the angle between the normal to the slip plane and the tensile
axis and X is the angle which the slip direction makes with the tensile axis, is therefore smaller than the corresponding value for aluminum. Schmid's law c
y
crss
cos 4> cos X
(4-1)
shows that the relatively low (cos(/> cos X) value is one reason (Sec. 4.2.2.2) for the larger hardness number and Hall-Petch coefficient (see also Sec. 4.4.9.2) of magnesium compared to aluminum, but it is also one reason why polycrystalline magnesium and its alloys do not develop macroscopic yielding and large stress concentrations at
124
4 Magnesium-Based Alloys
grain boundaries, instead. Pure magnesium and conventionally cast magnesium alloys thus have a tendency to embrittlement due to intergranular failure, and also localized transcrystalline fracture either along twinned regions, in particular upon compression, or along basal {0001} planes for very large grains (Emley, 1966). Prismatic {1010} <1120> slip planes are not active in pure magnesium (Fig. 4-5). The addition of lithium and indium can activate such slip planes when using conventional casting methods (Bach, 1969; Karney and Sachs, 1928; Hauser et al, 1958). Activation of prismatic {1010}<1120> slip renders magnesium alloys more ductile at lower temperatures without the need to recur to the required grain refinement as is obtainable by rapid solidification processing (RSP, Sec. 4.4.9.2). Unfortunately, however, many such conventional alloys are associated with extremely low amounts of solid solution-hardening and a low coldhardening response, as well as with poor corrosion resistance. Unless sufficient grain refinement was introduced to reduce back stresses and to enhance crystalline rotation via grain boundary gliding, however, deformation of structural magnesium alloys at ambient temperatures always shows a tendency to twinning (Sec. 4.2.2.3). Obviously the grain refinement of magnesium alloys offers a larger potential to improve their strength and ductility compared to aluminum alloys, and an interesting question is why this advantage has so far been left relatively uninvestigated, considering the advantages of this for other materials. Furthermore, what degree of grain refinement substitutes for the missing independent slip systems? Grain refinement is certainly one of the most important factors for the further development of magnesium alloys. However, conventional methods have merely the potential to in-
duce the grain size that assures the required ductility transition. Moreover, the comparison of modulus, operative slip normal vector, and resultant PeierlsNabarro stresses on the one hand with hardness and boundary strengthening capacity on the other indicates quite clearly that the mechanical behavior of magnesium is dictated by preferred dislocation reactions, as it is for other materials. Such dislocation reactions can take over the control of mechanical properties otherwise dictated by Peierls stresses (Sec. 4.2.2.2). If this was not the case, grain refinement of magnesium would do the reverse and magnesium would be softer than pure aluminum without showing better ductility. 4.2.2.2 Rate-Controlling Dislocation Reaction The scope of dislocation slip planes in the h.c.p. lattice covers basal, prismatic, and first order pyramidal slip planes with the Burgers vector b=lA a <1120>, and first and second order pyramidal systems with the Burgers vector lA (c + a) <1123> (Fig. 4-5). The V$ a <1120> vector is the shortest Burgers vector in the h.c.p. lattice (Fig. 4-6). The rate-controlling mode for plastic deformation is therefore dictated by the lattice friction experienced by the structure of the core of the lA a <1120> screw dislocation. Screw dislocations move slower than edge dislocations, since they are always present on two crystallographic planes and form sessile dislocation jogs upon propagation and intersection with other screws (Chen and Pond, 1952; Johnston and Gilman, 1957). The structure of the dislocation core tells with certainty whether energy is to be gained upon dissociation of perfect dislocation movements into imperfect or partial dislocation movements which leave a trace in the transla-
4.2 Characteristics of Magnesium a)
J
125
' [0001] (0110) prismatic
(2110) prismatic
repeat cell for ^-surface in b)
Figure 4-6. Illustration of the agreement between the intrinsic fault vector tl as derived from (a) linear elasticity and (b) multi-body potential calculations, here for the y-surface repeat cell (V^XllSO) <1100> of the basal plane of pure magnesium with a Burgers vector b (Vitek and Igarshi, 1991). Black discs: (0001) atoms; dotted discs: e.g. (0002) atoms; crosshatched discs then (0002) atoms, accordingly.
tional lattice, i.e. a stacking fault (SF) related to a given slip plane, and whether sessile SF components are involved. The operative slip mode of h.c.p. metals is confined either to the basal plane (simple metals like magnesium or beryllium) or to the prismatic plane (transition metals, see Table 4-2). The c/a ratio does not control the activation of prismatic slip as stated in some misleading literature (cf. Table 4-2).
Any comparison of the mechanical properties of magnesium-based alloys has therefore to be made on the basis of the structure of the core of the lA a <1120> screw dislocation. The basal dissociation of the V* a <1120> screw dislocation (Fig. 4-6) follows (Nabarro, 1967; Hirth and Lothe, 1968) (4-2) V3a<1120> -> y6a<1010> + Vfe
126
4 Magnesium-Based Alloys
Table 4-2. Glide modes in hexagonal close-packed
metals. h.c.p.element Cd Zn Co Mg Be Re Tc Tl Ru Os Hf Zr Ti
Principal glide mode a B B B B B B/P _ B/P P — P P P
Secondary Axial ratio c/a glide mode a nj.nj.p
n 2 ,p
n2
n2)p P,n 2 — — — — B,n2 n l 5 B, n 2 nl9
B,
n2
1.886 1.856 1.624 1.624 1.568 1.615 — 1.598 1.582 — 1.581 1.593 1.588
a B: basal plane; P: prismatic plane; H1 and U2 are first and second order pyramidal planes, respectively.
and corresponds to an SF energy yh of the order of 10 mJ m~ 2 for magnesium (Vitek and Igarshi, 1991). The prismatic SF energy yp is about seven times larger, which is why no cross slip occurs at ambient temperatures (see Sec. 4.2.2.1). For comparison, the SF energy y±11 of unalloyed aluminum is of the order of 200 mJ m ~ 2 (Dieter, 1986). The low basal stacking fault energy thus results in a low number of operative slip modes (Sec. 4.2.2.1), but it also allows for a larger dissociation of the ratecontrolling lA a <1120> basal screw than of the corresponding screw in pure aluminum. The dissociated basal screw is a moderately harder, transcrystalline obstacle to the motion of further dislocations, thus causing the pile-up of more dislocations in pure magnesium than in a hypothetical aluminum crystal without activated cross slip. This is the second factor contributing to the moderately higher values of hardness and ky, and to intergranular embrittlement of pure magnesium, though the Peierls stresses are much lower than in pure aluminum. The cold working capacity of pure
magnesium is thus slightly better than that of pure aluminum. For reasons of crystal symmetry it was postulated (Vitek and Igarshi, 1991), however, that any further dissociation, such as by intersection with other dissociated dislocations, cannot introduce sessile edge components into the basal plane. The basal stacking faults may always be glissile, as is evidenced by the relatively low work hardening response of magnesium alloys compared to aluminum alloys. This subject has been reviewed by one of the authors (Hehmann, 1995 b) the discussion there taking into account more recent empirical results (see Sec. 4.4.2.4). 4.2.2.3 Effect of Grain Size and Thermal Activation on Deformation Modes
Magnesium shows a marked ductility transition when twin formation is suppressed by grain refinement and by pyramidal {10ll}<1120> slip at temperatures above 225 °C (Fig. 4-7; Emley, 1966). Twinning is an athermal shear deformation process involving mirrored movement of several atomic layers at small fault vectors. Twin phenomena on second order pyramidal planes in magnesium thus reduce the large Peierls stresses associated with the large Vs (c + a)<1123> Burgers vector Fig. 4-5). Any increase in the susceptibility to thermal activation would therefore render slip deformation modes more competitive and would decrease the likelihood of twinning. There are three factors which increase the susceptibility of magnesium to thermally activated slip: (i) grain refinement, which reduces boundary back stresses so allowing easier accommodation of (a) the overlap or void displacement between two adjacent grains (i.e., grain boundary sliding and/or rotation), and (b) the twinned transcrystalline volume (Sec. 4.4),
127
4.2 Characteristics of Magnesium
(ii) alloying elements with low melting points (e.g. lithium), and (iii) increased temperature. Twinning is frequently observed at high deformation rates such as under cyclic loading of conventionally processed MgAl-based alloys (Attari et al., 1990). Grain refinement of such alloys results in significantly improved resistance to fatigue (Das et al., 1992). Grain refinement below a grain size of 8 jim for pure magnesium shifts the ductility transition down to room temperature (Fig. 4-7). The critical resolved shear stress (crss) for pyramidal slip in magnesium was found to peak at around 100°C, followed by decreasing pyramidal crss towards higher temperatures (Stohr and Poirier,
Table 4-3. Nomenclature used for magnesium alloy designations0; ASTM system. Letter A B C D E F G H K L
Element
Letter
Element
Aluminum Bismuth Copper Cadmium Rare earths (misch metala) Iron Magnesium Thorium Zirconium Lithium
M N P Q R
Manganese Nickel Lead Silver Chromium
S T W Y Z
Silicon Tin Yttrium Antimony Zinc
Examples15: AZ91: Mg-9%Al-l%Zn AZ91D: another specification, same general composition WE43: Mg-4%Y-3% rare earths a
Misch metal (MM): here 50% Ce, 20% La, 20% Nd, and 10% Pr; b unless stated otherwise, all compositions are given in mass%; c no designation has yet been established for calcium additions to magnesium.
50
100 150 200 250 300 Test temperature in °C
350
-200
-100 0 100 200 Test temperature in °C
Figure 4-7. (a) Effect of grain size on the temperature of the ductility transition and on the shape of the transition curve, (b) Effect of a gradual decrease of the grain size on the transition temperature. Both for pure magnesium. (Emley, 1966).
1972). Alloying with aluminum and zinc slightly reduces the brittle-to-ductile transition temperature [e.g. for AZM to 212°C (for magnesium alloy designations see Table 4-3); Emley, 1966]. When both suitable alloying and grain size refinement come together, however, magnesium alloys such as RS AZ91 (RS: rapid solidification) with a grain size of about 1 jim become superplastic even at room temperature. Elongations of more than 1000% were reported (Sec. 4.4; Lahaie et al., 1992). Novel micro-grained magnesium alloy compositions obtained by rapid solidification contain many temperature stable dispersoids which impede grain boundary sliding at lower temperatures. Their grain size is between 0.1 and 1 jim and they show a 400 % superplasticity in the temperature range of 275-300 °C (Das etal., 1992), which is better than for any of the other light alloys.
128
4 Magnesium-Based Alloys
4.2.3 Corrosion Behavior of Magnesium
Magnesium is a very unnoble metal. In an NaCl solution magnesium has the lowest electrical potential of all structural metals, i.e. magnesium is anodic to any other structural metal. Due to its standard potential of FNHE= - 2.37 V (NHE: normal hydrogen electrode), it oxidizes readily. In the absence of water a thin film of MgO is formed. MgO reacts with water to form Mg(OH) 2 , which has a pH value of 10.4. The resulting surface layer renders magnesium amenable for structural applications as it protects the metal from further oxidation, similar to the situation for other unnoble structural materials like, for example, titanium or aluminum. Unprotected magnesium may be more resistant to atmospheric corrosion than mild steel (Whitby, 1963). The Pourbaix diagram (Pourbaix, 1974) depicts the stability of magnesium in aqueous environments (Fig. 4-8). Magnesium is stable over the whole pH range at potentials VNHE= — 2.37 V or below, depending on the concentration of magnesium ions present in solution, as indicated in Fig. 4-8. At higher potential values and at pH values between 8.4 and 11.5, magnesium corrodes rapidly forming a layer of Mg(OH) 2 . This layer inhibits oxygen reduction at the surface and provides protection against corrosion. Compared to the respective surface layers on titanium or aluminum, however, the passivation layer on magnesium is less stable and more permeable. In addition to thermodynamic criteria, kinetic effects are important for the corrosion of magnesium. Figure 4-9 shows a typical polarization curve for AM20 in an Mg(OH) 2 solution. At potentials ranging from VNHE = — 1.4 to +1.4 V the alloy is passive because of kinetic suppression of the evolution of oxygen and
MgO 2 ?
o -2 -4 -6 )
c
Mg"
o
Mg(OH)2 (passive)
-2
@
h— •i
Mg+?
—'
s
•^ ^
Mg (passive)
-3 -2
8
10
12
14
16
pH-Value Figure 4-8. Pourbaix diagram showing the occurrence of various transformation products of the MgO film in aqueous solution as a function of applied potential and pH value at 25 °C (Pourbaix, 1974). At potentials below the broken line "a" hydrogen is evolved, above "b" oxygen is evolved, between "a" and "b" no dissociation of water occurs. Circled numbers relate to the boundaries of regions of different reaction products. The numbers beside the lines represent the logarithmic values of magnesium ion concentration in aqueous solution.
hydrogen. However, the addition of a few chloride ions leads to pitting corrosion. The corrosion behavior of magnesium alloys depends strongly on the metallurgy, microstructure, and alloying elements, for example. Certain impurities decrease the corrosion resistance of magnesium dramatically. The corrosion rate of magnesium in 3 % NaCl solutions, as well as in salt spray tests, can be improved by reducing impurities such as iron, copper, and nickel to levels below those shown in
129
4.2 Characteristics of Magnesium
Fig. 4-10 (Hanawalt et al., 1942; Reichek etal, 1985; Olsen, 1992). Reichek etal. (1985) reported the nickel tolerance limit as increasing with cooling rate, while the iron tolerance limit was observed to correspond to a constant value of 3.2% of the manganese content, independent of the casting conditions. The presence of aluminum reduces the tolerance limit of iron by about one order of magnitude, the correspondingly low solid solubilities are assumed (Froats et al., 1987) to lead to the formation of small intermetallic compounds which act as microgalvanic cells where the cathodic reduction of water readily occurs. Robinson and George (1954) suggested that manganese reduces the corrosion rate of magnesium by surrounding iron particles, because the galvanic potential between magnesium and manganese is lower than between magnesium and iron. Iron is detrimental to the corrosion resistance of magnesium only if it is present as inclusions. Homogeneously distributed iron atoms in magnesium lead to passivation, so improving the corrosion resistance (Akavipat et al., 1985; Sec. 4.4.8). The influence of other alloying elements on the corrosion resistance is less dramatic (Fig. 4-11; Hanawalt et al., 1942; see also Table 4-8). However, silver and calcium also have a strongly detrimental effect on the corrosion resistance of magnesium. Additions of rare earth (RE) elements and yttrium improve the corrosion resistance of magnesium. The standard potential of rare earth elements is similar to the value for magnesium (Biihler, 1990), resulting in low galvanic activity for the resulting RE compounds. In RS-processed Mg-RE alloys the corrosion improvement was considered to be due to RE enrichment in the surface (Krishnamurthy et al.5 1988). The effect of heat treatment and grain size on the corrosion rate of HP
-
1
0
1
2
NHE-Potential in V
Figure 4-9. Polarization curve for AM20 magnesium alloy in 0.003 M Mg(OH)2 solution, aerated, T= 30 °C, polarization rate = 1.2 V/h. _ 8
100 1.000 Iron, Nickel and Copper in \ig/g
Figure 4-10. Influence of nickel, iron, and copper impurities on the corrosion rate of die-cast AZ91 in the standard salt spray test (Olsen, 1992; Reichek et al., 1985).
Na,Si,Pb,Sn.Mn,Al_ 1
2
3
-jsJrTPb.Sn.Al U
Alloying element in %
Figure 4-11. Influence of alloying elements on the corrosion rate of magnesium as determined by immersion in 3 % NaCl solution (Hanawalt et al., 1942).
130
4 Magnesium-Based Alloys
Table 4-4. Typical corrosion rates of "high purity" AZ91E versus heat treatment3 and grain size (Froats et al., 1987). Alloy AZ91E
Untreated Degassed and grain refined Untreated Degassed and grain refined
Grain size
Corrosion rate in mm/year
Mn, Fe content b
in jim
Mn
Fe
146 78 160 73
0.23 0.26 0.33 0.35
0.008 0.008 0.004 0.004
F
T4
T6
T5
0.64 2.2 0.35 0.72
4 1.7 3 0.82
0.15 0.12 0.22 0.1
0.12 0.12 0.12 0.1
a
T4: 16 h 410°C/quenched; T5: 4 h 215 °C; T6: 16 h 410°C/quenched + 4 h 215 °C; F: as cast; analyzed Mn content.
AZ91E (HP: high purity) is shown in Table 4-4. The corrosion rate of AZ91E decreases with decreasing grain size in the heattreated condition, but increases in the ascast condition. The corrosion behavior of magnesium varies with the environment. Atmospheres free from salts do not attack magnesium alloys and a gray protective film is formed at the surface. In rural and industrial atmospheres magnesium is moderately attacked. Chloride contaminants at the surface destroy the surface film. Magnesium chloride has a high solubility in water and an acidic nature which causes the accumulation of anodic sites, so leading to microgalvanic couples and severe local attack. The corrosion behavior of magnesium in stagnant fresh water is similar to that in the atmosphere. The typical corrosion rate of AZ91 alloys in pure water at 100 °C is 0.25-0.5 mm/year, for example (Bothwell, 1967). Salt and in particular chlorides increase the corrosion rate. However, fluorides form insoluble magnesium fluoride and are not appreciably corrosive. Chromates, vanadates, phosphates, and some others are protective film forming and so retard corrosive attack. Magnesium exhibits good stability against various chemicals. It is resistant to
b
fraction of
dilute alkaline solutions. However, acidic solutions, except hydrofluoric acid and H 2 CrO 4 , attack magnesium rapidly. Magnesium is known as the sacrificial anode for cathodic protection. Direct contact of magnesium to electrically conductive materials results in macro-galvanic coupling, and can be suppressed by isolation and elimination of moisture. Surface modification of magnesium alloys can protect them from corrosive attack. Useful anorganic and/or organic coatings are given by Murray and Hillis (1990), for example. Novel nonequilibrium methods to improve the corrosion behavior of magnesium are rapid solidification processing and surface alloying (see Sec. 4.4.8). 4.2.4 Further Properties of Engineering Relevance
The electrical resistivity of polycrystalline magnesium Q at 20 °C is 4.46 x 10 ~8 Qm (Kirk-Othmer, 1981), which is about three times that of copper. The resistivity ratio QcJQa for a magnesium single crystal is 0.83 (Roberts, 1960). The thermal expansion of magnesium is relatively large compared to most other structural materials. The linear coefficient of thermal expansion (CTE) increases
131
4.2 Characteristics of Magnesium
from 25 x 10 6 /K at ambient temperature to 30 x 10" 6 /K at 527 °C. This is roughly 10% more than the respective values for aluminum. There is only a small difference between the thermal expansion coefficients of a single crystal of magnesium along the a-axis and the c-axis. The expansivity ratio CTEa/CTEfc is 0.96. On going from the melting temperature to ambient, solid magnesium has a linear contraction of 1.9%. Magnesium alloys have a high damping capacity. The specific damping capacity is defined as the ratio of the energy absorbed per cycle to the total energy per cycle of a material in vibration. The maximum value for cast, pure magnesium is reduced by alloying additions, strengthening treatments, and working. Low-alloyed, cast magnesium has a better specific damping capacity than cast iron (Emley, 1966). Some data on the specific damping capacity of magnesium are given in Table 4-5. The low specific heat capacity of magnesium may cause some pre-solidification problems, e.g. in gravity die casting, but offers the opportunity to reduce the shot time in any casting process. In general, the diversity of casting conditions for magnesium is restricted compared to aluminum. Magnesium is the most easily machinable of all the structural metals. Except for drilling, the energy required to machine magnesium-based alloys is nearly half of that for aluminum-based alloys (Table 46). Relatively small chip dimensions facilitate computer-controlled machining of magnesium and magnesium alloys. Machining of magnesium alloys instead of aluminum alloys increases the hard metal tool lifetime by a factor of 5-10. Machined magnesium surfaces are usually smooth without the need to recur to lubrication at tool speeds below 1000 m/min. The use of lubricators is recommended, however, in
Table 4-5. Specific damping capacity of some magnesium alloys (Housh et al., 1990). Alloy
Temper a
Specific damping capacity in % 14MPa
35MPa
AM60 AS21 AS41 AZ91 K1A
F F F F F
13 33 13 5 49
52 60 44 29 66
EZ33 ZE41
T5 T5
5 2
22 2
A1356
T6
0.5
Cast iron
F
5
1.2 17
a
F: as fabricated; T5: artificially aged only; T6: solution heat treated and artificially aged.
Table 4-6. Machinability of different materials (Kato et al., 1992; Magnesium Die Casting Manual, 1993). Material
Cutting energy in kW/ms Lathing
Mg alloys Al alloys Cu alloys Ti alloys Mild steel Cast iron SUS 304(18-8)a
Drilling
Milling
1.95 1.95 1.95 3.1 1.95 4.2 8.1-12.8 6.1-10 8.1-12.8 15.3 13.9 13.9 13.9-43.1 12.8-32.8 13.9=32.8 8.9-17.8 12.8-20.3 7.5-13.9 16.4-17.8 13.9-15.3 17.8-18.9
SUS 304(18-8) = AISI 304 (stainless steel).
order to avoid ignition and to explore the machinability of magnesium alloys at tool speeds as high as 7500 m/min. Subsequent recycling procedures form part of established alloy transformation schemes (see Sec. 4.6). 4.2.5 Summary of the Advantages and Limitations
The basic advantages and limitations of magnesium and its alloys are summarized
132
4 Magnesium-Based Alloys
Table 4-7. Basic advantages and limitations of magnesium and its alloys. Advantages
Limitations
Lightest structural material (alloy densities £ = 1400 to 1900 kg/m3) High casting production rate due to low heat content Compatible with mild steel as material for crucibles, pipes, etc. Faster to machine with less power than other structural materials High specific strength
Low number of established alloy families
High damping capacity Less corrosive that iron Fully recyclable
in Table 4-7. The reader will find in the following Sees. 4.3-4.5 that numerous developments are under way in order to remove such limitations and to explore the advantages of magnesium more effectively.
4.3 Magnesium Alloys and Their Applications 4.3.1 Equilibrium Alloying
The development and use of conventional magnesium alloys started at the beginning of the second decade of this century and peaked in the late 1930s. A second peak in magnesium alloy development was in the 1950s (Serve, 1962). In the last 40 years or so magnesium alloy development has been less exciting than the development of steel or aluminum, for example. This may be partly due to the advent of plastics and fiber-reinforced plastics as materials competing with the conventional magnesium alloys that had been developed until recently. Conventional magnesium alloys are strengthened by solid solution hardening
Small number of wrought alloys and products Highly reactive melt with some of the alloying elements Low stiffness compared to other metallic materials High thermal expansion compared to other metallic materials In general, I/M produced alloys are more corrosive than most aluminum alloys Widespread prejudice of easy burning and corroding No existing secondary metal market
and/or by precipitation hardening. Extensive solid solutions form if the atomic radii of the solute and the solvent do not differ by more than about 15% (see, e.g., Carapella, 1945). Figure 4-12 shows the atomic radii of the elements in relation to magnesium and their maximum solubilities in solid magnesium. The greater the difference in atomic size, the more restricted is the solubility provided that additional factors governing extensive solid solutions including type of crystal structure, equivalent valencies, and electrochemical factors are favorable (Carapella, 1945). There is a tendency for the solid solubility to become more restricted as the valencies become more unequal. Provided atomic size is favorable the most important restriction for large solid solubilities is magnesium's strong electropositivity (Carapella, 1945). The maximum solubility in magnesium is observed for elements of group IIB of the periodic table, namely h.c.p. zinc and cadmium. Only cadmium forms a continuous solid solution with magnesium at elevated temperatures (>253 °C). For efficient solid solution hardening, the difference in atomic radii of solute and solvent should be as large as possible. Usu-
4.3 Magnesium Alloys and Their Applications
133
100 Figure 4-12. Atomic radii of the elements (bottom) and their maximum solubility in solid magnesium (top). The horizontal solid line represents the atomic radius of magnesium, the horizontal dashed lines a deviation of + 15% from the atomic radius of magnesium.
134
4 Magnesium-Based Alloys
ally a significant solid solution hardening effect is expected if the solid solubility of the alloying element is higher than about 0.5 at.%, provided the difference in atomic radii is sufficient. Precipitation hardening requires a reduction in the solid solubility with decreasing temperature, and the resultant precipitates have to be stable at the service temperature. Magnesium is relatively electropositive, resulting in the formation of intermetallic compounds with most alloying elements. The stability of the compound tends to increase with increasing electronegativity of the alloying element (Carapella, 1945). The precipitation sequence usually starts with small coherent particles which grow and/or coarsen and loose coherency at elevated temperatures (Polmear, 1992). Some important alloying elements are very expensive, such as the rare earth elements, yttrium, and silver, limiting their use to specific applications. The most commonly used alloying elements are aluminum, zinc, manganese, and zirconium. Rare earth metals and yttrium are used for elevated temperature applications. Table 4-8 gives a short overview of the effect of several common alloying elements in magnesium. The three major alloy systems at present are magnesium-aluminum alloys, magnesium-zinc alloys, and magnesium-rare earth metal alloys. At present magnesium does not have a rich alloy spectrum like aluminum or zinc. Fresh stimulation for alloy development is needed. More recent approaches include micro-alloying to redesign existing magnesium alloys with trace additions of surface active elements like Ca, Sr, Ba, or Sb, Sn, Pb, and Bi. For elevated temperatures, new magnesium alloys have been developed using rare earth element additions.
4.3.1.1 Magnesium Casting Alloys Magnesium alloys are mainly used for automobile components, housings, and electric components, which are produced by a die-casting process. The general advantages of the die-casting process are: high productivity, high precision, good casting surface, good and fine solidification structure, and production of thinwalled and complicated-shape cast components. Die casting of magnesium can easily result in gas porosity due to splashing during mold filling. The die-casting process is not very suitable for manufacturing thickwalled castings. Die casting can take advantage of certain characteristic properties of magnesium, such as its lightness, high specific strength, good machinability, high damping capacity, good castability, and low melting temperature and melting energy. Typical die-casting alloys are given in Table 4-9. Their physical properties are listed in Table 4-10. Magnesium die-casting alloys are based on the Mg-Al system. Alloying with aluminum results in strength and good castability for these alloys. An aluminum content of more than 3 % is necessary for easy die casting. Zinc is also used for strengthening magnesium alloys, but it tends to cause hot cracking at concentrations above 2%. Manganese produces Al-Mn-Fe compounds which (a) sink into the sludge of the melt, thus reducing the iron content and (b) refine microstructure. Fe, Ni, and Cu are harmful to the corrosion resistance of Mg-Al alloys (Sec. 4.2.3). The Fe content and the ratio of Fe/ Mn are strictly controlled, and this improves the corrosion resistance of magnesium cast materials remarkably. These alloys are known as "high purity" (HP) alloys. Figure 4-13 illustrates the dramatic effect of these impurities on the corrosion
4.3 Magnesium Alloys and Their Applications
135
Table 4-8. General effects of several alloying elements on magnesium. Alloying element
Melting and casting behavior
Ag Al
Improves castability, tendency to microporosity
Be
Significantly reduces oxidation of melt surface at very low concentrations (< 30 ppm), leads to coarse grains Effective grain refining effect, slight suppression of oxidation of the molten metal System with easily forming metallic glasses, improves castability Magnesium hardly reacts with mild steel crucibles
Ca Cu Fe Li
Mechanical and technological properties
Corrosion behavior I/M produced
Improves elevated temperature tensile and creep properties in the presence of rare earths Solid solution hardener, precipitation hardening at low temperatures (< 120 °C)
Detrimental influence on corrosion behavior
Improves creep properties
Detrimental influence on corrosion behavior
Increases evaporation and burning behavior, melting only in protected and sealed furnaces Control of Fe content by precipitating Fe-Mn compound, refinement of precipitates System with easily forming metallic glasses
Solid solution hardener at ambient temperatures, reduces density, enhances ductility
Rare earths
Improve castability, reduce microporosity
Si
Th
Decreases castability, forms stable silicide compounds with many other alloying elements, compatible with Al, Zn, and Ag, weak grain refiner Suppresses microporosity
Solid solution and precipitation hardening at ambient and elevated temperatures; improve elevated temperature tensile and creep properties Improves creep properties
Y
Grain refining effect
Zn
Increases fluidity of the melt, weak grain refiner, tendency to microscopy
Zr
Most effective grain refiner, incompatible with Si, Al, and Mn, removes Fe, Al, and Si from the melt
Mn Ni
Increases creep resistivity
Improves elevated temperature tensile and creep properties, improves ductility, most efficient alloying element Improves elevated temperature tensile and creep properties Precipitation hardening, improves strength at ambient temperatures, tendency to brittleness and hot shortness unless Zr refined Improves ambient temperature tensile properties slightly
Minor influence
Detrimental influence on corrosion behavior, limitation necessary Detrimental influence on corrosion behavior, limitation necessary Decreases corrosion properties strongly, coating to protect from humidity is necessary Improves corrosion behavior due to iron control effect Detrimental influence on corrosion behavior, limitation necessary Improve corrosion behavior
Detrimental influence
Improves corrosion behavior Minor influence, sufficient Zn content compensates for the detrimental effect of Cu
136
4 Magnesium-Based Alloys
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rate in a standard salt spray test. This limitation on the impurities has improved the corrosion resistance of AZ91 by a factor of about 30 compared to the "old" AZ91. The corrosion rate of AZ91HP is lower than that of the aluminum alloy A l Si9Cu3. Table 4-10 shows that the specific gravity and the energy of fusion per unit volume of the magnesium alloys are very low. These are the advantages of magnesium alloys, but they do cause some difficulty in casting design. The casting temperature of the melt in magnesium die casting is between 650 and 680 °C for AZ91, 670 and 685 °C for AM60, and 665 and 680 °C for AS41. The energy needed to heat magnesium AZ91 from ambient to 50 °C above the melting temperature is 480kcal/dm 3 ; this value is 77% of that of the ADC 12 (Al-Sil2CuFe) aluminum die-casting alloy. Magnesium die-casting alloys have traditionally been melted by using some type of flux protection. A typical flux salt mixture is, e.g. DOW230 (KC1 55 vol.%, MgCl2 34 vol.%, CaCl 2 9 vol.%, CaF 2
137
4.3 Magnesium Alloys and Their Applications
Table 4-10. Physical properties of magnesium die-casting alloys. Property Specific gravity in g/cm3 Solidification range in °C Specific heat in J/kg K Thermal conductivity in W/m K Thermal expansion coefficient in 10" 6 /K Heat of fusion in kJ/kg a
Aluminum die-casting alloy;
b1
AM60
AS41
AZ91
ADC12a
ZDC2 b
1.78 610-510
1.78 620-562 1.01 67 26
1.81 596-568 0.98 52 26 372.4
2.8 582-516 0.96 100 21 389.1
6.6 387-381 0.42 113 27 100.4
61 26
zinc die-casting alloy.
2vol.%). Sulfur powder is also used on top of the magnesium melt to suppress ignition. Today, SF6-containing gas mixtures are used to protect the magnesium melt. This obviates the inclusion of flux in the cast materials, and it also greatly improves the corrosion resistance of die-cast materials. Typical protective gas mixtures are (Magnesium Die Casting Manual, 1993): (1) AZ91D 650-705 °C Air+ 0.2% SF 6 (all in vol.%) (2) AZ91D 650-705 °C 75% Air+ 25% CO 2 + 0.2% SF 6 (3) AZ91D 705-760 °C 50% Air+ 50% CO 2 + 0.3% SF 6 SF 6 concentrations above 0.5% do not result in significant further improvements in the protection of the magnesium melt but tend to attack the steel of the die, die chamber, crucible, etc. The addition of small quantities of beryllium to magnesium alloys is very effective in restraining burning. Usually a quantity of 5-10 jig/g Be is added to magnesium die-cast alloys. The employed die-casting methods use hot chamber and cold chamber technology. In the hot chamber machine the piston is situated inside the melt (Fig. 4-14). The advantages of hot chamber die casting are: high productivity, low temperature of melt and long life of die, and good protection of the melt. However, its main disadvantages
are expensive machinery, and complicated and expensive maintenance. Hot chamber die casting is primarily used for relatively small magnesium alloy components. In the cold chamber machine the piston is separated from the melting furnace (Fig. 4-15). The corresponding advantages and disadvantages are almost the opposite of those of the hot chamber machine. Furthermore, a special system is needed to transport the melt from the furnace to the shot cylinder in order to protect the magnesium melt from oxidation and ignition. Several pumping systems have been proposed (Magnesium Manual, 1991). New die-casting technologies such as vacuum die casting (Pennington, 1993) and squeeze die casting have recently been applied to magnesium alloys (Chadwick Hydraulic Clamp
Plunger Shot Cylinder Molten Metal-
Figure 4-14. Hot chamber die-casting machine.
138
4 Magnesium-Based Alloys Hydraulic Clamp
Figure 4-15. Cold chamber die-casting machine.
and Bloyce, 1992); the latter mainly for producing magnesium matrix composites (Sec. 4.5). To explain the microstructure of AZbased magnesium alloys the Mg-Al binary phase diagram is shown in Fig. 4-16 (Massalski et al., 1990). The eutectic temperature is 437 °C and the maximum solu-
bility of aluminum at this temperature is 12.7%. The solubility decreases with decreasing temperature. At 200 °C, for example, h.c.p. magnesium can dissolve 3.2% aluminum. The microstructure of as-cast AZ91 consists of a-magnesium solid solution and the interdendritic Mg 17 Al 12 intermetallic phase. Die-cast magnesium alloys have a zinc content of 1-3% and a manganese content of less than 1 %, both having only a minor effect on the microstructure. Figure 4-17 shows the microstructure of the die-cast AZ91 alloy. For comparison the microstructures of plaster-cast and sand-cast AZ91 alloys are shown in Fig. 418 and Fig. 4-19, respectively. The relatively high cooling rates of the die-casting process result in a non-equilibrium mi-
Atomic Percent Magnesium 10 700
600-1
500-
400-
a
g
300-
200-
100 20
Al
30
40
50
60
70
80
Weight P e r c e n t Magnesium
Figure 4-16. Aluminum-magnesium binary alloy phase diagram (Massalski et al., 1990).
90
100
Mg
4.3 Magnesium Alloys and Their Applications
139
(a) (a)
|_Mgs.s. + MgtTAIii]
10
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(b) (b)
(C)
(c)
Discontinuous phar.e]
Figure 4-17. Microstructure of die-cast AZ91 alloy: (a) as cast, (b) after T4 treatment, (c) after T6 treatment (for treatment details see Table 4-4).
Figure 4-18. Microstructure of plaster-cast AZ91 alloy: a) as-cast, (b) after T4 treatment, (c) after T6 treatment (courtesy of M. Pekgiileryiiz, ITM).
crostructure. The as-cast condition shows segregations of Al-poor (Mg dendrites) and Al-rich regions of magnesium solid solution (solidified remainder of the liquid). After T6 (Table 4-4) heat treatment, a discontinuous precipitation grows from the Al-rich areas at the grain boundaries. The
white phase on the grain boundary of slowly solidified castings (Figs. 4-18 a and 4-19a) is the intermetallic phase Mg 17 Al 12 . An Al-enriched region has been formed around the white Mg 17 Al 12 phase. During further cooling, lamellar Mg 17 Al 12 precipitates are formed in the area supersaturat-
140
4 Magnesium-Based Alloys
curs inside the grains and the precipitates are aligned with the {0001} and the {1120} planes of the magnesium-based solid solution. The discontinuous precipitation starts at grain boundaries and dominates at ageing temperatures below 300 °C. Early studies on the ageing of Mg-Al alloys showed (Fox and Lardner, 1943; Leontis and Nelson, 1951):
(a)
SQfim (b) ,
(c)
Figure 4-19. Microstructure of sand-cast AZ91 alloy: (a) as cast, (b) after T4 treatment, (c) after T6 treatment (courtesy of M. Pekgiileryuz, ITM).
ed with Al. A major proportion of the Mg 17 Al 12 segregation and of the fine Mg 17 Al 12 precipitates dissolves on subsequent solution heat treatment (Figs. 4-18 b and 4-19 b). Both continuous and discontinuous Widmanstatten-like precipitates are observed after ageing (Figs. 4-18 c and 4-19c). The continuous precipitation oc-
(1) improved strength by suppressing discontinuous precipitation and accelerating continuous precipitation; (2) increased ratio of continuous over discontinuous precipitation with increasing cooling rate after artificial ageing; (3) improved strength with increasing refinement of precipitates; (4) that the conditions for the finer precipitate were lower ageing temperature, shorter ageing time, and lower aluminum content; (5) no confirmation of any evidence of coherency between the matrix and the above precipitates. The mechanical properties of die-cast magnesium alloys are shown in Table 4-11. These values are for the standard die-cast condition, and the actual strength of magnesium cast materials may differ from these data. AZ91 has high strength at room temperature due to the added zinc but at the expense of a reduced elongationto-fracture (EA). The AM series is virtually free of zinc (Table 4-9) and was mainly developed to enhance ductility at the expense of strength and hardness. The AM20 alloy shows EA values in the range of 10 %. Recent investigations showed that elongations of 20 % are possible with AM-type alloys. Reducing the aluminum content slightly increases the corrosion rate of AM-type alloys (Fig. 4-13). Zn-containing magnesium alloys require low heating rates at 300 to 350 °C in
141
4.3 Magnesium Alloys and Their Applications
Table 4-11. Mechanical properties of magnesium die-casting alloys. Property Tensile strength in MPa 0.2% proof stress in MPa Elongation in % Brinell hardness number Young's modulus in GPa Fatigue strength in MPa (5 x 107 cycles) a
Aluminum die-casting alloy;
AM20
AM50
AM60
AS41
AZ91
ADC12a
ZDC2 b
160-210 90-120 8 - 12 4 0 - 55
180-220 110-140 5- 9 50- 65
190-230 120-150 4- 8 5 5 - 70 43 4 9 - 69
190-230 120-150 3- 6 60- 90
196-255 147-167 0.5- 3 6 5 - 68 44 49- 69
220-320 180-220 1- 3 70-100 76 60- 90
265-304
b
4 9 - 69
10 82 91 69
zinc die-casting alloy.
Table 4-12. Influence of different heat treatments on mechanical properties of a permanent mold cast magnesium AZ91 alloy. Heat treatmenta
Tensile strength
Proof stress
Elongation
Homogenization treatment
AZ91-F AZ91-T4 AZ91-T5 AZ91-T6
> 160 MPa > 240 MPa > 160 MPa >240 MPa
> 70 MPa > 70 MPa > 80 MPa > 110 MPa
>7% >2% >3%
410-420 °C 16-24 h 410-420 °C 16-24 h
Artificial ageing treatment
170°Cforl6h 170°Cforl6h
215°Cfor4h 215°Cfor4h
a
F: as fabricated; T4: heat treated and aged at room temperature to a stable condition, but more than 50 % of alloying elements remain in solution; T5: artificially aged only; T6: solution heat treated and artificially aged.
order to avoid local remelting of Zn-enriched microstructural zones. The standard conditions for the heat treatment of AZ91 are given in Table 4-12. The temperature dependence of the mechanical properties of die-cast AZ91 are shown in Fig. 4-20 (Magnesium Manual, 1992). The tensile strength of die-cast AZ91 decreases above 120°C due to insufficient thermal stability of the Mg 17 Al 12 phase. The creep properties of AZ91 and AS41 are shown in Fig. 4-21 (Mercer, 1990). An alloy for service temperatures of slightly more than 200 °C is one of the main targets of present research and development. As described above, the strength of AZ91 drops at temperatures of more than 120 °C. The first such alloy was AS41, which is now approved by the die-casting standards of the ASTM. AS-type magne-
sium alloys were already in use in the VW "beetle". The creep properties can be improved further (Fig. 4-21) by reducing the aluminum content (AS21), but this alloy is more difficult to cast. Production of AS41 requires a relatively fast solidification process such as by high pressure die casting in order to avoid the formation of the coarse Chinese script morphology of the Mg2Si phase, which reduces the fluidity and the resulting elongation. Pekgiileryuz and Avedesian (1992) have shown that small additions of calcium modify the structure of AS41 in sand casting and refine the Mg2Si-phase. Improvements in the UTS and elongation were reported upon the addition of 0.1 % Ca to AS41. The same effect was observed in gravity die-cast AS41. Figure 4-22 shows the unmodified and the Ca-modified microstructures of AS41.
4 Magnesium-Based Alloys
Figure 4-20. Ultimate tensile strengths of selected magnesium die-cast alloys as a function of temperature (Ninomiya et al., 1995) [AC33 = MgAlCa33 (Mg3% Al-3% Ca-0.6% Mn)]. 300
350
400
450
500
Temperature K
200
Figure 4-21. 0.1 % creep strength after 100 h of some magnesium die-cast alloys as a function of temperature (Mercer, 1990). Data for ZC63 from Magnesium Elektron Ltd. (courtesy of X King, MEL). Data for 380A1 are given for comparison. 100
150
200
250
300
Temperature in °C
Figure 4-22. Micro structure of (a) gravity die-cast AS41, and (b) gravity die-cast 0.13% Ca-modified AS41, as cast, etched. The Chinese script morphology of Mg2Si disappears upon small additions of Ca.
4.3 Magnesium Alloys and Their Applications
Another approach was taken by Magnesium Elektron Ltd. (MEL) using an M g Zn-based alloy. The development of alloy ZC63 resulted in enhanced creep properties and good castability. This alloy is free from microshrinkage and is weldable (Gwynne et al., 1988). Despite the detrimental effect of minor Cu additions on the corrosion resistance of magnesium alloys, fully heat-treated ZC63 shows better corrosion resistance than AZ91C in the salt spray test (Sec. 4.2.3; see also Hanawalt et al., 1942). The alloy ZC63 can be used in sand casting, gravity die casting, and precision casting. AE42 (Mg-4% A l - 2 % misch metal (MM) - 0 . 6 % Mn) has been developed for the same application field in North America (Mercer, 1990; see Sec. 4.3.1.2). Many attempts to use this alloy for automobile components have been carried out and some practical applications have resulted. AE42 requires fast solidification such as in high pressure die casting in order to achieve good mechanical properties. Ninomiya et al. (1995) have studied Mg-Al-Ca alloys and developed the alloy MgAlCa33 (Mg-3 % Al-3 % C a 0.6%Mn). MgAlCa33 has good die-castability and good high temperature mechanical properties (Fig. 4-20). The MgAlCa33 alloy shows improved resistance to creep over AE42 due to a ratio of Ca/Al greater 0.8, resulting in substantial replacement of the Mg 17 Al 12 dispersion by thermally more stable Mg2Ca-based precipitates (see Sec. 4.4.3.2). 4.3.1.2 Magnesium Alloys for Elevated Temperature Application
A number of existing magnesium alloys can be used at elevated temperatures, including AS- and ZC-type alloys if the service temperature is limited to about 150-
143
200 °C. Magnesium alloys applicable at service temperatures above 200 °C, however, are essentially magnesium-rareearth-based alloys. Magnesium forms solid solutions with most of the rare earth elements (Fig. 4-12). The magnesium rich corners of the corresponding binary systems are simple eutectic. This also applies when cheaper Cebased or Nd-based misch metals are added to magnesium. The eutectic binary phase diagrams underlie the good casting characteristics of most of the magnesium-rare earth alloys. The corresponding low melting point eutectics form a grain boundary network which tends to suppress microporosity (Polmear, 1992) and to accommodate the primary grains of the magnesium solid solution (Fig. 4-23). Subsequent ageing results in transcrystalline precipitation strengthening. The creep properties of the magnesium-rare earth alloys are excellent due to grain boundary precipitation reducing sliding of the grain boundaries. The precipitation behavior of some of the magnesium-rare earth alloys such as Mg-Nd is not fully clear. Coherent GP (Guinier-Preston) zones form at temperatures below 200 °C (see, e.g., Polmear, 1992) and may be of plate-like morphology in the Mg-Nd system. In Mg-Ag-Nd alloys rod-like or ellipsoidal coherent precipitates were observed, and in the M g - Y - N d system such GP zones have not yet been confirmed (Polmear, 1992). The GP zones transform into coherent hexagonal p" plates of D0 19 superlattice structure and of composition corresponding to the Mg 3 Nd phase. The latter precipitates were not found in Mg-Ag-Nd alloys. At temperatures above 200 °C most Mg-RE alloys have semicoherent p' precipitates such as Mg 3 Nd. Overageing results in magnesium-rich, incoherent precipitates. The structure and composition
144
4 Magnesium-Based Alloys
Figure 4-23. Microstructure of the magnesium alloys WE 54 and binary Mg-3.1 % Sm in the as-cast condition and after annealing for homogenization (T4); cross sections, etched: (a) WE54, as-cast condition, permanent mold casting; (b) WE54, T4; (c) Mg-3.1 % Sm, as-cast condition, permanent mold casting; and (d) Mg-3.1 % Sm, T4.
of corresponding ternary and quaternary systems are currently being investigated (cf. Ahmed et al., 1992). The first improvements in the elevated temperature tensile strength of Mg-Ce alloys were reported by Haughton and Prytherch in 1937 and by Beck in 1940. Beck shows earlier (1926) diagrams by Menking. Subsequent work by Leontis (1949) showed that La, Ce, and Nd improve the elevated temperature properties in that order. A big step forward was made when Sauerwald (1947) discovered the grain refining effect of zirconium. Zr can be used together with alloying elements
like Ag, Zn, rare earth metals, and Th. The alloying elements Al and Mn form stable intermetallic compounds with Zr, so limiting the use of Zr as a grain refiner. Sauerwald's findings led to the development of EK-type magnesium alloys (Table 4-13). It was found that magnesium casting alloys with Nd misch metal gave better properties than with Ce misch metal. Small amounts of Zn improve the room temperature properties. Rare earth metal and Zn additions have opposite effects on magnesium alloys. Rare earth additions to M g Zn alloys improve their castability and creep resistance but reduce their tensile
4.3 Magnesium Alloys and Their Applications
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properties. A compromise has to be made depending on the intended application. For this purpose the alloys ZE41 and EZ33 were developed. ZE41 is a high strength alloy for use at temperatures up to 200 °C, whereas EZ33 is a more creep resistant alloy for use at temperatures up to 250 °C (Unsworth, 1989). Zn additions are detrimental to the corrosion behavior of Mg-RE alloys. In 1959 Payne and Bailey discovered that silver additions significantly improve the age hardening response of Mg-RE alloys. Following these results alloys like QE22, QH21, and EQ21 (Table 4-13) were developed. This particular alloy development again confirmed Leontis' observation that Nd-based misch metal additions result in better properties than Ce-based misch metal additions. The addition of Ag significantly reduces the corrosion resistance of Mg-RE alloys. Silver-containing Mg-RE alloys show good tensile properties (Fig. 4-24) at ambient temperatures. At temperatures up to 200 °C alloys QE22 and EQ21 show yield strengths of 175 MPa or more (Fig. 4-24 b). The yield strength of QE22 decreases rapidly at temperatures of more than 200 °C, whereas EQ21 exhibits excellent strength properties up to 250 °C. The QE alloys offer good resistance to fatigue (Fig. 4-25). The fatigue strength of QE22 at 200 °C exceeds 80 MPa. The creep strength of the QE alloys is comparable to the creep strength of EZ33 (Fig. 4-26). The above QE-type alloys have elevated temperature tensile properties and creep resistance that are nearly as good as for the thorium-containing magnesium alloys. Thorium is the most effective alloying element with magnesium for giving elevated temperature properties. Zinc- and silvercontaining magnesium-thorium alloys have been developed. The low radioactivi-
4 Magnesium-Based Alloys 50
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i
100
200
300
400
300
200
300
200
300
400
Temperature in °C
Temperature in °C
100
i
100
400
Temperature in °C Figure 4-24. Temperature dependence of tensile mechanical properties of some magnesium-rare earth alloys (T6: WE54, WE43, QE22, EQ21; T5: ZE41, EZ33; high pressure die cast: AE42); (a) ultimate tensile strength, (b) yield strength, and (c) total strain (courtesy of X King, MEL).
ty of Th, however, restricts the application of Mg-Th-based alloys and excludes them from future applications. A very important discovery was the beneficial effect of yttrium additions to magnesium, leading to a family of Nd- and Y-containing alloys with high strength and creep properties at elevated temperatures (Drits et al, 1979; Unsworth, 1989). The WE-type alloys have excellent corrosion resistance, similar to that of the common aluminum casting alloys. One example is the sand-cast alloy WE54. WE54 contains 4.75-5.5% yttrium, 1.5-2% neodymium, 0.4-1.0% zirconium, and up to 2% heavy rare earth elements (Table 4-13). Yttrium is usually added in the form of yttrium misch metal, containing about 75 % yttrium with the remainder as heavy rare earth metals. This misch metal does not affect the precipitation sequence and only slightly reduces the strength compared to the addition of pure yttrium (Gwynne et al., 1988).
4.3 Magnesium Alloys and Their Applications 140
140
a)
I
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Room Temperature
120
120
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80
60
1 20
40
20
1E+05
1E+06
1E+07
1E+08
1E+05
1E+06
1E+07
1E+08
Number of Cycles
Number of Cycles
Figure 4-25. Fatigue strength (unnotched) of some magnesium-rare earth alloys at (a) room temperature and (b) 200 °C (courtesy of J. King, MEL).
200
a) 200C 150
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100
Time in h
1.000
100
Time in h
Figure 4-26. Creep strength of some magnesium-rare earth alloys (T6: WE54, WE43, EQ21; T5: EZ33; high pressure die cast: AE42) at (a) 200 °C and (b) 250 °C (courtesy of J. King, MEL).
148
4 Magnesium-Based Alloys
Y and Nd additions increase the volume fraction of eutectic phases in the as-cast condition, resulting in detrimental, intragranular segregation. Subsequent solution treatment is employed to dissolve the eutectic phase, but a small volume fraction of Nd-based eutectic remains (Fig. 4-23). The precipitation sequence (see, e.g. Ahmed et al., 1992) starts at between 170 and 200 °C with the formation of hexagonal P", which has a D0 19 superlattice structure parallel to the {1010} magnesium matrix planes, followed by the formation of body centered orthorhombic P' at between 200 and 250 °C. These precipitates also have an orientation relationship to the magnesium matrix. Above 300 °C, equilibrium f.c.c. Pprecipitates form heterogeneously within the grains and at grain boundaries. The tensile strength, fatigue strength, and creep strength of WE54 are excellent at ambient and elevated temperatures (Figs. 4-24 to 4-26), and are comparable to aluminum cast alloys. After an exposure of more than 1000 hours to slightly elevated temperatures, the ductility of WE54 at ambient decreases below the specified 2%. At a service temperature of 150°C, p" precipitates form and coexist with the already present p'-phase (Ahmed et al., 1992). The embrittlement increases with the volume fraction of P", which depends on the difference in the solid solubility at the ageing temperature and the service temperature. The alloy WE43 contains slightly less Nd and Y, resulting in somewhat lower strength properties but satisfying ductility (Gwynne et al., 1988). The mechanical properties of WE43 are excellent compared to other magnesium alloys (Figs. 424 to 4-26). WE43 is used in racing cars and for aeronautical applications such as the transmission casing of the McDonnell Douglas MD500 helicopter (King and Wardlow, 1993; Fig. 4-27).
Figure 4-27. WE43 transmission casing of the McDonnell Douglas MD500 helicopter (courtesy of X King, MEL).
Experimental magnesium alloys containing Sm, Gd, and heavy rare earth metals have been investigated more recently. The solubility of the REs in Mg except Eu and Yb - increases to about 10at.% for the heavy rare earth metals (Fig. 4-12). Extraordinary mechanical properties can be achieved with Gd and Dy additions, but more than 10% Gd or Dy is required to obtain these properties. Figure 4-28 shows, for example, the tensile data for the alloy Mg-9.4%Gd4.8%Y-0.6%Mn (Drits et al., 1979). The high rare earth metal price is one obstacle to more substantial consideration from the user-side. An interesting field for future investigations is magnesium-samarium alloys (see Neite et al., 1993). Sm has a maximum solubility of 5.7% in magnesium and offers good solid solution and precipitation hardening. More than 10% of the heavy rare earth elements from Gd to Yb is required to achieve an equivalent hardening response (see Fig. 4-12). The tensile strength of the binary Mg-6%Sm alloy in the T6 condition is comparable to that of WE43, but its yield strength is lower (Fig. 4-28; Drits et al., 1985).
149
4.3 Magnesium Alloys and Their Applications 500
500
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200
300
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Temperature in °C Temperature in °C Figure 4-28. Temperature dependence of tensile mechanical properties of an Mg - Sm and an Mg - Gd alloy (T6). Data for WE43 and AE42 are shown for comparison: (a) ultimate tensile strength and (b) yield strength.
The most frequently used alloying element for magnesium is aluminum. The presence of RE metal in combination with aluminum results in the formation of very stable RE aluminides, thus removing the RE metal from the magnesium solid solution. Therefore this combination of alloying elements cannot be used, e.g. in sand casting or permanent mold gravity die casting where the cooling rates are fairly low. Experiments have shown that the cooling rate obtained by die casting is sufficiently fast to substantially suppress the formation of aluminides. The development of several AE-type alloys at Dow (Mercer, 1990) took advantage of the cooling rate obtainable by die casting and resulted in finely dispersed aluminides and good creep properties up to 300 °C (Fig. 4-21). The magnesium alloy AE42X1 has a slightly better creep strength than that of AS21, a tensile strength similar to AS41, a salt-
water corrosion resistance similar to AZ91D, and its castability is only slightly inferior to that of AS41A (Baker, 1992). Another attempt was made by Beer et al. (1992) to develop an elevated temperature Mg-Si alloy for piston applications. The investigated alloy contained 10-50 vol. % of the intermetallic line compound Mg2Si, which melts congruently at 1085 °C. The low thermal expansion and the high Young's modulus of Mg2Si can compensate for some of the drawbacks of magnesium. A large volume fraction of dispersed Mg2Si is thus very useful and desired. The high liquidus temperature promotes the formation of primary Mg2Si dendrites. RE-containing inoculants were shown to lead to a homogeneous distribution of Mg2Si crystals with an octahedral morphology (Fig. 4-29). A further addition of aluminum strengthens the Mg matrix and improves the technological properties at
150
4 Magnesium-Based Alloys
OQ QQ GQ CQ PQ 0Q PQ PQ CQ QQ PQ
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elevated temperatures. The Young's modulus and the thermal expansion coefficient of a magnesium alloy with about 10% Si are comparable to common cast aluminum alloys. This inexpensive alloy conception also provides a high specific stiffness and strength up to temperatures as high as 300 °C (Beer et al., 1993).
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4.3.2 Wrought Magnesium Alloys It is necessary to develop wrought magnesium products such as rolled sheet or plate, extrusions (bars, rods, shapes, and tubings), and forgings in order to achieve more substantial structural applications of magnesium alloys. These wrought forms have the advantages of lower cost, higher strength and ductility, as well as better versatility of the mechanical properties than in the cast form. Progress has been made in the extrusion technology of magnesium, for example (Kittilsen and Pinfold, 1992). The compositions of typical wrought magnesium alloys are shown in Table 4-14. For wrought magnesium alloys produced by the various rapid solidification methods, the reader is referred to Sec. 4.4.
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4.3 Magnesium Alloys and Their Applications
4.3.2.1 Processing As a result of the limited deformability of h.c.p. Mg (Sec. 4.2.2), hot working is usually employed as the initial forming process in order to accomplish large reductions in the cross section without crack formation in the corresponding preform. Hot working reduces the number of roll passes and reheating intervals required to transform the cast billet into its final dimensions prior to finishing. Hot rolling is performed between 300 and 450 °C, depending on the alloy composition. The reduction for one hot rolling pass ranges from 10 to 30%, and the number of reheating intervals to complete hardening during hot working ranges from 0 to 5, depending on the alloy used. Typical sequences of hot rolling of AZ31 are as follows: (i) Slab rolling: starting temperatures range from 425 to 450 °C, reduction per pass of 10-20% resulting in a total reduction of 90-95%. (ii) Coil rolling: starting temperatures range from 350 to 440 °C, reduction per pass of 5-20 % resulting in a total reduction of 25-50%, 3 reheating intervals.
(iii) Final rolling to obtain the best combination of mechanical properties, dimensional accuracy, and stability as well as surface quality is as follows: starting temperatures below 250 °C, reduction per pass of 5 % resulting in a total reduction of 15-25 % followed by annealing at a temperature of about 135°C. A summary of typical rolling schedules for magnesium wrought alloys is shown in Table 4-15. The extrusion of magnesium is similar to that of other metals. The applicable force for standard extrusion presses in commercial production range from 1200 to 15 000 tons (1200-150001) {Magnesium Handbook, 1975). The summary of typical extrusion schedules for commercially extruded alloys is shown in Table 4-16. It is noted that extrusion yields a wide range of mechanical properties depending on the working temperature employed. Hot extrusion is performed at about 400 °C, at which temperature the extruded alloys are already recrystallized. Cold extrusion is usually done at temperatures below 300 °C where many microstructural features of cold work, such as a high density of dislo-
Table 4-15. Summary of typical rolling schedules for magnesium wrought alloys. Alloy
Operation
Temperature
Reduction per heat cycle
Reduction per pass
AZ31
Slab breakdown rolling Coil rundown rolling Finishing step Slab breakdown rolling Coil rundown rolling Finishing step Slab breakdown rolling Coil rundown rolling Finishing step Slab breakdown rolling Coil rundown rolling Finishing step
425-450°C 350-440 °C <250°C 450-500 °C 350-450 °C room temperature 480 °C 480 °C 150-250 °C 480-500 °C 375-425 °C room temperature
90-95% 25-50% 15-25%
10-20% 5-20% 5% 10-30% 10-40% 5%
Ml
HK31
HM21
40-60%
5 - 7% 30-40% 26%
3-10% 3 - 5%
152
4 Magnesium-Based Alloys
Table 4-16. Summary of typical extrusion schedules for magnesium wrought alloys. Alloy
Billet temperature
Container temperature
Extrusion rate
AZ31 AZ61 AZ80 Ml
371-400 °C 371-400 °C 360-400 °C 416-438 °C
232-316°C 232-288 °C 232-288 °C 380-388 °C
15- 40 7 - 20 4- 7 20-100
Table 4-17. Temperature dependence of mechanical properties for extruded magnesium wrought alloys. Alloy
AZ31
Ml Pure Mg
Operation Tensile temperature strength in °C in MPa 380 400 430 250 430 130 170 210 280 330
300 270 245 275 235 245 240 235 235 230
Yield strength in MPa
Elongation in %
230 185 150 185 135 145 145 140 140 135
21 16 14 7 6 4.5 4 4 4 4
cations or a twinning microstructure, are retained inside the material. The better mechanical properties of extruded products are obtained by extrusion at a temperature somewhat below the recrystallization temperature (Table 4-17). The tensile strength of AZ31 increases from 245 to 300 MPa, the yield strength from 150 to 230 MPa, and the elongation from 14 to 2 1 % on decreasing the extrusion temperature from 430 to 380 °C. However, with pure magnesium metal the extrusion temperature hardly affects the mechanical properties. 4.3.2.2 Mechanical Properties at Room Temperature
The room temperature mechanical properties of wrought magnesium alloys are shown in Table 4-18 for rolled sheet and plate (Magnesium Elektron Ltd., 1961; Dow Chemical Company, 1964 a, b, 1983 a, b; Smithells, 1976), in Table 4-19 for extruded bars, rods, and shapes (Dow
Table 4-18. Typical mechanical properties at room temperature of magnesium alloys in the form of rolled sheet and plate. Alloy
AZ31B
MIA HK31A HM21 ZE10A HZ11A ZK10A ZK30A ZM21A a
Temper3
O H24 H26 O H24 T8 O H24 O H24
O
Tensile strength
Compressive yield strength in MPa
Elongation
in MPa
Tensile yield strength in MPa
250-255 255-290 260-275 232 220-230 255-270 230-255 215-230 235-260 235 260 263 270 233
145-150 145-220 170-205 100 115-140 200-215 160-185 110-160 165-200 125 170 178 185 131
75-110 85-180 105-165
17-21 14-19 10-16 6 17-23 9-14 10-12 18-23 8-12 6 12 10 8 13
90- 95 150-170 125-160 105-110 160-180
154 154
in %
O: softest temper (annealed, recrystallized); H24: strain hardened then annealed (medium, 50 % hard); H26: as H24, 75 % hard; T8: solution heat treated, cold worked and artificially aged at elevated temperature.
153
4.3 Magnesium Alloys and Their Applications
Table 4-19. Typical mechanical properties at room temperature of magnesium alloys in the form of extruded bars, rods and shapes. Alloy
AZ10A AZ31B or C AZ61 AZ80A MIA HM31A ZC71A ZH11A ZK10A ZK30A ZK60A ZN21A
Tempera
F F F F T5 F T5 T6 F F F F T5 F
Tensile strength
Compressive yield strength in MPa
Elongation
in MPa
Tensile yield strength in MPa
204-240 260 310-315 330-340 345-380 225 305 356 263 293 309 330-340 360-365 255
145-150 195-200 215-230 240-250 260-275 180 270 330 147 208 239 250-260 295-305 162
70- 75 95-105 130-145
10 14-15 15-17 9-12 6- 8 12 10 5 18 13 18 9-14 11-12 11
215-240 125 160-185 141
213 160-230 215-250
in %
F: as fabricated; T5: artificially aged; T6: solution heat treated and artificially aged.
Table 4-20. Typical mechanical properties at room temperature of magnesium alloys in the form of extruded tubing. Alloy
AZ10A AZ31B AZ61 MIA ZK10A ZK60A a
Tempera
F F F F F F T5
Compressive yield strength in MPa
Elongation
in MPa
Tensile yield strength in MPa
230 250 285 240 278 325 340
145 165 165 145 193 240 270
70 85 110
8 12 14 9 7 13 12
Tensile strength
175 180
in %
Codes as in Table 4-19.
Chemical Company, 1965; Smithells, 1976; ASM, 1979; Busk, 1988), in Table 4-20 for extruded tubes (Dow Chemical Company, 1965; Smithells, 1976; ASM, 1979), and in Table 4-21 for forgings (Dow Chemical Company, 1962b, 1967; ASTM, 1984). Figure 4-30 shows typical stress-strain curves for extruded AZ31 under compressive and tensile loading at ambient temperature. Note that at lower temperatures ex-
trusion and rolling tend to align the basal planes parallel to the direction of extrusion and rolling. If the grain size is large enough, as for I/M processed magnesium alloys, twinning occurs readily upon compression parallel to the basal plane (Emley, 1966). The resulting, overall drawback of I/M processed magnesium alloys is the lower compressive yield strength compared to the tensile yield strength, and this
154
4 Magnesium-Based Alloys
Table 4-21. Typical mechanical properties at room temperature of magnesium alloys used for forgings. Alloy
AZ31B AZ61B AZ80A
EK31A HK31A HM21A ZH11A HK30A ZK60A a
Temper3
Compressive yield strength in MPa
Elongation
in MPa
Tensile yield strength in MPa
260 195 315 345 345 290 260 235 232 309 305 325
195 180 215 235 250 195 195 150 147 224 205 270
85 115 170 195 185 155 140 110
9 12 8 6 5 7 21 9 13 8 16 11
Tensile strength
F F F T5 T6 T6 F T5 F F T5 16
193 195 170
in %
Codes as for Table 4-19.
is particularly evident for wrought I/M magnesium alloys due to their texture. Only a grain size below about 10 jim can increase the ratio of the compressive-to-tensile yield strength towards unity and above (see Sees. 4.2.2 and 4.4.2.2). 300 As-extruded AZ31B-F 250 tensile
compressive
0.6
0.8
1.0
1.2
Strain
Figure 4-30. Typical true stress-strain curves of asextruded AZ31B-F alloys tested in tension and compression at room temperature.
The most important commercial wrought magnesium alloys are AZ31B or AZ31C, which differ in the allowable impurity content. AZ31's good combination of strength and ductility is mainly a result of controlled work hardening retained upon rolling, extrusion, or forging, and upon intentional or incidental annealing during cooling down to room temperature. The addition of aluminum and zinc to magnesium leads to solid solution hardening and grain refinement. I/M AZ61A is one of the extruded and forged magnesium alloys with both excellent strength and ductility due to a higher content of aluminum than for I/M AZ31 alloys, for example. The higher aluminum content, however, tends to result in cracking upon rolling, and this alloy is rarely supplied in sheet form. Extruded ZK60 has been developed as a new and promising wrought magnesium alloy. Big advantages over Mg-Al-based extrusions and forgings are the tensile and compressive yield strengths of more than 280 and 210 MPa, respectively, as obtained after ageing, and a relatively large fracture strain of more than 10%. On a
155
4.3 Magnesium Alloys and Their Applications
4.3.2.3 Mechanical Properties at Elevated Temperatures Rolled and extruded magnesium alloys have been developed for elevated temperature applications. It was found that the addition of rare earth metal or thorium reduced grain growth, the tendency to recrystallize, and grain boundary deformability. The first magnesium-based sheet was made from the alloy HK31A-H24. This was the cold-worked equivalent of the high-temperature casting alloy of the same composition. The temperature dependence of tensile strength, yield strength, and fracture strain of magnesium alloy sheets is shown in Figs, 4-31, 4-32, and 4-33 (Fenn and Lock-
1
300
strength-to-weight basis this high strength is equal to 420 MPa for an aluminumbased wrought 7075 or 7475 alloy in the aged condition.
Rolled magnesium -O-AZ31B/S-O -^-AZ31B/S-H24 -*-HK31A/S-O —^-HK31A/S-H24 -Q-HM21A/S-T8 -•-HM21A/P-T8
+ ft
s. ^
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I .22
I4
"§ 100 • -¥-ZE10A/S-H24 -*^-ZH11A/S-H24 - f f l - ZK30A/S-H24 0
200
V\\
400 600 Temperature, K
\ 800
Figure 4-32. Temperature dependence of the tensile yield strength of magnesium alloy sheets.
500
CO Q.
AZ31B/S-0 AZ31B/S-H24 HK31A/S-O HK31A/S-H24 HM21A/S-T8 HM21A/P-T8 ZE10A/S-O ZE10A/S-H24 ZH11A/S-H24 ZK30A/S-H24
400
2
160 140 120
£ 100 O O)
"55 o
80
Rolled magnesium — o - AZ31B/S-0 - • - AZ31B/S-H24 —o— HK31A/S-0 — * - HK31A/S-H24 - n - HM21A/S-T8 — • - HM21A/P-T8 — v - ZE10A/S-O - ¥ - ZE10A/S-H24 - N - ZH11A/S-H24 —Ba— ZK30A/S-H24
60 40 20
. Rolled magnesium 200
400
0 600
800
Temperature, K Figure 4-31. Temperature dependence of the tensile strength of magnesium alloy sheets.
200
400
600
800
Temperature, K
Figure 4-33. Temperature dependence of the elongation of magnesium alloy sheets.
156
4 Magnesium-Based Alloys
Tensile strength 500
400
I £
300
I to
"55 Q>
200
100
Extruded magnesium
100
200
Figure 4-34. Temperature dependence of the tensile strength of extruded magnesium alloys.
300
400
500
600
Temperature, K
500 Tensile yield strength' -•-4-o-^-•-a-^-
400 CO Q.
§ 300 V> 2 w
100
AZ31B-F AZ61A-F AZ80A-F HM31A-F HM31A-T5 M1A-F ZK30A-F ZK60A-T5
Extruded magnesium
80 -
60
200
Elongation - • - AZ31B-F - • - AZ61A-F -o- AZ80A-F -^- HM31A-F - » - HM31A-T5 - a - M1A-F -^- ZK30A-F - * - ZK60A-T5
o LU
100
40
20
100
200
300
400
500
600
Temperature, K
Figure 4-35. Temperature dependence of the tensile yield strength of extruded magnesium alloys.
100
200
300
400
500
600
Temperature, K
Figure 4-36. Temperature dependence of the elongation of extruded magnesium alloys.
157
4.3 Magnesium Alloys and Their Applications
wood, 1960; Fenn, 1961; Dow Chemical Company, 1964 a, c, 1967, 1983 a; Smithells, 1976), respectively. The temperature dependence of tensile strength, yield strength, and fracture strain of extruded I/M magnesium alloys is shown in Figs. 4-34, 4-35 and 4-36 (Dow Chemical Company, 1962a, 1967; Smithells, 1976; ASM, 1979), respectively. With increasing temperature the tensile strength decreases, whereas the fracture strain increases. Large fracture strains of more than 100% are found at temperatures above 350 °C depending on the alloy composition. Embrittlement at low temperatures was not observed in wrought magnesium alloys. It is well known that thorium-containing alloys such as HM31XA show good high
temperature properties up to about 370 °C. Recently, some extruded magnesium-rare earth metal alloys have been developed which are equally strong at higher temperatures. The details of these magnesium alloy systems are described in Sec. 4.3.1.2. 4.3.2.4 Magnesium-Lithium Alloys Magnesium-lithium alloys (Fig. 4-37) are the lightest structural metals with excellent deformability and superplasticity in the eutectic composition range (Fig. 4-38; Metenier et al., 1990; Gonzalez-Doncel etal., 1990; Higashi and Wolfenstine, 1991; Higashi etal., 1992, 1993; Taleff etal., 1992). The as-cast (oc + p) microstructure of a eutectic Mg-8.5% Li alloy
Atomic Percent Lithium 20 30
40
50
60
700
70
80
100
95
90
v650°C 600-
500-
400-
(Li)
(us) ^
300-
' i
200-
L
\
180.6°C:
600-^
1X,692°C
100-
(Kg)
.5
580(3 0
Mg
10
20
• i
• • i
5 30
10 40
\ N \ \ 15 50
Weight P e r c e n t
60
Lithium
Figure 4-37. Mg-Li binary alloy phase diagram (Massalski et al. 1990).
70
80
90
100
Li
158
4 Magnesium-Based Alloys
1000
hep
bec
hep+bec
800
600
Elongat
P
623 K
Superplasticity region ~~
\
e =4x10- 4 S" t _
200
O
D-
°
J
Figure 4-38. Typical elongation behavior of Mg-Li binary alloys.
n 4
8
12
16
20
24
Li content, wt%
shows Widmannstatten-type a-platelets in a p-matrix (Fig. 4-39). Extensive work at Batelle Memorial Institute (see Jackson et al., 1949; Jackson and Frost, 1967) based on early studies in the 1930s and 1940s (see, e.g. Grube etal., 1934; Dean and Anderson, 1943; Shamrai, 1947) resulted in the development of the p-b.c.c. magnesium alloy LA141A (Mg-14%Li-l %A1) with a density Q of 1.35 g/cm3 through the addition of Li (Q = 0.55 g/cm3). Subsequent studies were devoted to the ageing behavior, workability, and microstructure of Mg - Li based alloys (McDonald, 1968; Lee and
Jones, 1974; Hafner, 1976; Schurmannand Hansen, 1982; Hansen et al., 1986; Schiirmann and Engel, 1986; Schemme, 1993 a). Alloy LA141A was used for aeronautical applications. The corresponding engineering properties are summarized in Table 422. Ternary additions X to Mg-14 at.%Li such as Al, Ag, Cd, and Zn were reported (cf. Jackson et al., 1949; Jackson and Frost, 1967) to provide more effective age hardening response by the formation of the coherent LiaXd phase. Susceptibility to overageing is one reason for the limited use of this alloy. Recently, (oc +13) Mg-8.5-9.0%Li alloys were prepared (Gonzalez-Doncel et al., 1990) by cold rolling and subsequent lowtemperature press-bonding of the thin foils into laminates prior to superplastic deformation. The resulting microstruc-
Table 4-22. Basic properties of as-cast LAI41A at ambient temperature Property
Figure 4-39. A typical microstructure of the as-cast, eutectic Mg-8.5% Li alloy. The light etched phase is the magnesium-rich a-phase.
Specific gravity Tensile strength Proof stress Elongation Young's modulus
LA141A 1.35 g/cm3 130 MPa 103 MPa 12% 42.7 GPa
4.3 Magnesium Alloys and Their Applications
ture consisted of fully recrystallized, finegrained, a- and p-phases of volume fraction 30% and 70%, respectively (Fig. 440 a; Metenier et al., 1990; Taleff et al., 1992). The final grain size of the laminates depended strongly on the total reduction of the foil thickness during cold rolling. The strain-rate-sensitivity exponent m of Mg-9Li laminates of grain size 3-20 \xm was reported to exceed values of 0.5 when temperatures between 150 and 250 °C and strain rates of the order of 10 ~2 s""1 were employed. A tensile elongation of 460 % was obtained at 180°C using a strain rate of 3 x l 0 " 4 s ~ 1 . More recently (Taleff et al., 1992), a fine-grained Mg-9Li laminate of grain size 1.5 pm resulted in an m value of 0,5 and an elongation of 450% at a temperature of 100 °C, for example. Higashi and Wolfenstine (1991) reported superplastic behavior for an Mg-8.5 % Li alloy on applying true strain rates between 10~ 5 s" 1 and 10~ 3 s" 1 at temperatures ranging from 250 to 375 °C. The microstructure of this alloy consisted of an unrecrystallized oc (40vol.%) and P (60 vol.%) two-phase mixture (Fig. 4-40 b) obtained by warm rolling at 200 °C and using a 7:1 ratio for the total reduction of the corresponding cross section. True constant strain rate tests resulted in 300610% elongation, the latter obtained at 350 °C by using a strain rate of 4 x l 0 ~ 4 s " 1 . Under the same testing conditions, monitoring the changes in the microstructure and the m values as a function of superplastic strain revealed that the microstructure changed from an initially banded (a + P) microstructure into homogeneously distributed, equiaxed (a + P) grains of 10 jim in size underlying superplastic deformability (Fig. 4-41), and the value of m increased from 0.4 to 0.7. These changes suggested that the Mg-8.5% Li alloy underwent continuous dynamic re-
159
(a)
(b) Figure 4-40. Typical microstructure of Mg-Li alloys produced by two different processing routes: (a) Mg-9% Li alloy processed by foil metallurgy. The rolling direction is vertical, (b) Mg-8.5% Li alloy after warm rolling at 200 °C. The rolling direction is horizontal.
crystallization during superplastic deformation. A ternary Mg-Li-Y alloy was reported to show superplasticity at relatively high strain rates (Higashi et al., 1992). The top of Fig. 4-42 shows the flow stress values of Mg-8.5% Li and Mg-8.5% Li-1 % Y as a function of the strain rate at r=350°C. An m value of about 0.5 was obtained at strain rates near 10~ 4 s" 1 for the binary alloy, and at between 10 3 s i and 10~ 2 s~ 1 for the ternary alloy, which showed lower flow stresses than the binary
160
4 Magnesium-Based Alloys
10
10
Strain rate, s"1
Figure 4-42. Variation in flow stress and elongation for the Mg-8.5% Li alloy and the Mg-8.5% Li-1 % Y alloy tested at 350°C as a function of the strain rate.
Figure 4-41. The microstructure of the Mg-8.5% Li alloy as a function of strain for a strain rate of 4x 10" 3 s" 1 at 350°C: (a) 8 = 0.7, (b) e = 1.4, and (c) after failure. The tensile direction is horizontal.
alloy. The bottom of Fig. 4-42 shows the total elongation values of the Mg-Li and the Mg-Li-Y alloys as a function of the strain rate at r=350°C. Elongations
above 300% were recorded for the binary alloy at strain rates oflO~ 4 -10~ 3 s~ 1 and for the ternary alloy at 10~3-10~2 s"1. The maximum elongation of the binary Mg-Li alloy is about 600 % at a strain rate of 2xl0" 4 s~ 1 , and it is 400% for the ternary Mg-Li-Y alloy at a strain of 4xl0~ 3 s~ 1 . Despite this lower maximum elongation, the ternary alloy was reported to show a higher superplastic strain rate than the binary Mg-Li alloy. It should be noted that the grain size of the Mg-Li-Y alloy is smaller than for the binary Mg-Li alloy (Fig. 4-43). This may result from the addition of yttrium, similar to the effect of higher order additions to RSP Mg-Al-Zn alloys (see Sec. 4.4.3). The high strain rate superplasticity of the Mg-Li-Y alloys, which results from refinement of the microstructure, is of benefit for commercial forming applications (Higashi et al., 1992).
4.4 Rapid Solidification Processing of Wrought Magnesium Alloys
161
(a)
Figure 4-43. Typical microstructures of (a) Mg-8.5% Li and (b) Mg-8.5% L i - 1 % Y alloys after annealing at 350 °C for 5 h. The rolling direction is horizontal.
4.4 Rapid Solidification Processing of Wrought Magnesium Alloys 4.4.1 Synopsis Alloy conversion without a change in the state of aggregation tends to find applications in tinker and niche markets. Powder metallurgy (PM) of light metals is such a niche market due to the limitations PM imposes on the productivity and the resultant global balance of the synthesis (and consolidation) of advanced light metals (Hehmann, 1992 a). That is why the PM of light metals has never formed an explicit part of the global classification of the processing routes for alloy conversion, despite the notable progress made during the last decade. Alloy conversion alternatives for monolithic and precursor alloys are classified instead with respect to the cooling rate employed for solid-state formation from the liquid and vapor phases (Fig. 4-44). Within this classification, conventional ingot metallurgy (I/M) is considered to range from cooling rates of 10" 3 K/s to about 103 K/s, followed by rapid solidification processing (RSP) from the melt with cooling rates of about 103 K/s to about 109 K/s (including a substantial range of gas atomization methods) and the ultra-rapid solidification method from the vapor phase [sublimation or physical vapor deposition (PVD)] with cooling rates greater than
10 10 K/s (Bianchi, 1991; Lavernia etal., 1993). Cooling rate considerations alone, however, cannot account for the phenomena encountered during solidification, such as the very different microstructures obtained by using the same type of chill material, for example. Rapid solidification processing involves by definition a high-velocity propagation of the solidification front. The operative driving force for local mass transport at the growth front is dictated by the rate of latent heat extraction during the change in the state of aggregation. This requires either a short solidification path in order to control the required heat removal from outside the transforming volume via heat transfer and chill agents, or a very high undercooling in order to accom-. 10 J 10 2
ias Water EBED
10°
EBEP
10" 10-
-
—Stainless
10" 10"3
10° 10 3 106 109 10 12 10 15 Cooling Rate (K/s)
Figure 4-44. Global classification of gas-, water-, and chill-block quenching solidification routes (Bianchi, 1991). (EBED, electron beam evaporated deposit; EBEP, electron beam evaporated powder.)
4 Magnesium-Based Alloys
modate the corresponding latent heat internally, so allowing for a temperature rise during recalescence without sacrificing the desired structural effects. In practice, however, inoculants triggered by alloy constitution, contaminations, process parameters, and process-dependent artifacts render the undercooling of large volumes impossible and make short solidification paths inevitable. Fragmentation to establish short solidification paths is achieved by gas or liquid atomization methods to form a population of droplets, and also by confined fluid flow such as by chill-block splat and spinning methods or momentarily confined fluid flotation by laser remelting for epitaxial growth of the surface layers, all of which involve growth normals with a scale limited by less or not much more than 100 |im of material. While the scale of the growth normal is set by the processing method concerned, the major question focuses on the effective driving force for local mass transport. This driving force is uncontrollable when nucleation barriers impose heat flow back into the liquid, as in industrial PM processes established to date, and any modeling has found its limitations there. For a system like Mg-Sr, the models by Clyne et al. (1984) and Boettinger et al. (1986) predict a difference in undercooling, AT^, of 40 K for droplet diameters d ranging from 1200 to 12 jam, i.e., for a difference in length scale and nucleation probability of at least two orders of magnitude, which is in disagreement with practical observations (see Fig. 4-45; Hehmann, 1994 a). Classical nucleation theory, however, predicts a AI N range of 200 K for wetting angles 9 ranging from 150° to 30°. Obviously, nucleation efficiency is ratecontrolling, but the available models do not explicitly relate nucleation efficiency to nucleation probability and the length
(a) 4-r
Mg - 5.9 at.% Sr
321to
log [
162
0 = 90° Recalesced Planar Growth
0- ^ \ / 12jLtin -1- ^ " * " ' ' * * * ^ ^ -21200 jurf -3-4-
Maximum Dendritic
vmax ^ J ^ ^
Dendritic Growth '
llK^
1
0.0
|
0.25
0.50
1.0
0.75
Cross section x/d
(b)
t
3-
Mg - 3.0 at.<;
21-
d = 12 Jim Maximum Dendritic
0-1-2-3-
-40.0
Vel cit
°
—^ \
o
\
\ c 45° 60 \
0.25
y V max X /
\X
1
75° \ \
0.50
0.75
90° \ ' 1.0
Cross section x/d
Figure 4-45. Transition from segregation-free to dendritic growth during recalescence as a function of (a) droplet size (12-1200 jam), and (b) nucleation efficiency (wetting angle 30°-90°) for hypo-eutectic magnesium-strontium alloys using the models by Clyne et al. (1984) for predicting nucleation temperatures as a function of droplet size, followed by using Lipton et al.'s (1987) model for free dendritic growth in undercooled alloy melts combined with a simple heat balance (Boettinger et al., 1986). vRS, recalesced solidification velocity; x, velocity position over fragmented cross section; d, fragmented size (Hehmann, 1994 a).
scale-dependent field of thermal diffusion during recalescence. In order to circumvent this problem, an increasing number of authors (Saunders and Tsakiropoulos, 1988; Pan et al., 1989; Miiller et al., 1992) attribute their observed microstructural features to the convention
4.4 Rapid Solidification Processing of Wrought Magnesium Alloys
that nucleation efficiency is related to a value 6 of the order of 70°-90° or so. The physical significance of such an interpretation is to render the nucleation rate independent of the alloy constitution and thermal diffusion ahead of the advancing solidification front and dependent on process parameters and length scale instead, without involving the effective nucleation efficiency. This approach conflicts with classical nucleation theory, as indicated by the results in Fig. 4-45, and it cannot replace the required model to relate nucleation probability to nucleation efficiency (Adkins and Tsakiropoulos, 1991; Shao etal., 1993). Any progress in the hierarchization of the involved physical parameters should therefore involve a more relevant nucleation theory and a more effective development of advanced light alloys, in which magnesium will play a more significant role than is yet appreciated (Hehmann, 1992 a, 1995 a). The following review also includes details of powder handling, degassing and storage in view of their effect on engineering properties wherever possible. 4,4,2 Brief Summary 4.4.2.1 Atomization Versus Chill-Block Methods
The development and production of commercially available magnesium alloys made by rapid solidification (RS) is concentrated over two periods: era 1 from about 1950 to 1960 at the Dow Chemical Co. using gas atomization (Chisholm and Busk, 1953) and later on rotating cooling discs (Colby and Hershey, 1955; Chisholm and Hershey, 1956; Fig. 4-46 a); and era 2 with the development from 1984 to date that was spurred by the development at Allied Signal (Fig. 4-46 b), and pursued further by about 15 organizations evenly
163
distributed throughout Europe and America. More recently, RS magnesium alloys have attracted attention in Japan (see Sees. 4.4.7 and 4.4.8). The invention by the Dow Chemical Co. in the late forties of direct extrusion of magnesium-based powders without precompacting and the need for atmospheric protection opened new possibilities by using pre-alloyed powders. In their classical paper, Busk and Leontis (1950) reported improvements that could not be fully realized by conventional means, i.e. sufficient grain refinement, interference hardening, suppression of the tendency to stress corrosion cracking, and improved extrudability by using co-extrusion of blends of different alloy powders rather than ingots. The latter was related to solid-state interdiffusion, resulting in the precipitation of finely dispersed, insoluble Al3Zr compounds, leading to the best thermal stability and strength values of up to 430 MPa reported in that era (Leontis and Busk, 1953). Interference hardening is a microstructural formation process of second phase dispersions having a melting point which is higher than or close to the boiling point of magnesium. The high melting point of these phases "interferes", for example, with ingot processing of corresponding alloy compositions. Support beams made from the Dow alloys were used for the floor or built-in loading ramps of the C133 transport plane, involving a 40% increase in the compressive yield strength (CYS) (PM ZK60 over ingot-cast, ZK60), for example (Foerster and Johnson, 1958). Today, the pioneering work by Dow Chemical still serves as a treasure chest, as is evident from the more recent developments at Allied Signal and Pechiney/ Norsk-Hydro, all of which largely benefit from the interference hardening approach
4 Magnesium-Based Alloys
Figure 4-46. Schematic diagrams of (a) the disc atomization unit at Dow Chemical (Colby and Hershey, 1955; Chisholm and Hershey, 1956) and (b) the planar flow casting (PFC) process at Allied Signal (Das et al, 1988) for rapid solidification processing of magnesium-base alloys. The Dow process (a) of the 1950s involved a melting unit with furnace setting (33), melting pot (36), and tubular spigot (37), all heated by flames provided through openings (47) and (48). The diameter of the outlet of the spigot (38) was smaller than the bore (39) of the spigot. It was closed by push rod (42) for the valve (41). The atomizing disc (27) was operated between 2000 and 100 000 rpm and was propelled by a gas turbine (21) with inlet (20) and exhaust (50) via vertical spindle (23), all situated in the powder collection tank under its concave portion (32). Allied's PFC unit (b) includes a chill-block (14) onto which the molten puddle (2) is forced through nozzle lips (3) and (4) to produce ribbons of final cross section (5) before flying of the wheel (15). A scraper (7,12) reduces the air boundary layer (1) under the ribbon so improving heat extraction, while an induction unit (9, 10) controls the liquid temperature of the melt. The side shields help to stabilize the melt puddle (2) in conjunction with the gas inlet (8).
via the formation of precipitates such as A12Y, A13RE, and Al2Ca, stemming from the otherwise incompatible alloy constituents in corresponding magnesiumbase liquids and the I/M processed solid state (see Sec. 4.4.9.3). Dominating RSP methods in era 2 were chill-block methods such as melt spinning (MS) and twin-roller (TR) quenching (Das et al., 1988; Fig. 446). Chill-block RSP methods have the advantage of avoiding highly flammable and adherent fines which occur during gas atomization, and the need to filter, clean, and recycle large volumes of corresponding, inert atomizing gases. Furthermore,
chill-block methods provide more effective heat extraction, affording supersaturation and more uniform microstructures of the resultant component than is obtainable by gas atomization methods, since the latter involve per se some microstructural nonuniformity as a result of the omnipresent variations in scale of powder particles. Gas atomization, however, does not require ribbon communition and is thus a somewhat shorter PM route. Gas atomized powders also allow some freedom in the fabrication of composites due to more regular and round-shaped alloy particles versus those made from communited ribbons,
165
4.4 Rapid Solidification Processing of Wrought Magnesium Alloys
including requiring a lower shearing force and press capacity for interparticle welding. Oxide dispersions in the Mg bulk alloy are a general obstacle to achieving sufficient fracture toughness and resistance to corrosion when dispersed in bulk magnesium alloys, and this spurred some of the more recent efforts via atomizing and spray forming, for example. The latter have not reached, however, the status that was achieved by employing chill-block methods (Sees. 4.4.3 to 4.4.5). This subject was reviewed by Hehmann (1984), Hehmann and Jones (1986, 1987, 1990, 1993), Lewis etal. (1987), Jones et al. (1987), Froes etal. (1987, 1988), Das and Davis (1988), Chang et al. (1989), Jones (1989, 1993), and Bray (1990), and is here reviewed again with respect to the more recently evolving trends and highlights.
pilot-production scale, proprietary planar flow casting (PFC) process developed at Allied Signal. PFC is performed using a rectangular slot orifice as a casting nozzle in close proximity to the rotating chill wheel. A shroud allows a jet of inert gas to stabilize the shape of the molten pool of metal on the rotating chill block (Fig. 4-46; Das et al., 1988), so arriving at larger widths of the resultant ribbons than is possible by free jet casting or "melt spinning". The process enables a spinning operation of continuous ribbons of 25 ^im thickness and of an average width of 50 mm, followed by pulverization to sub-micrometer flakes which are then used in ordinary PM and mill processing. The process arrives finally at product forms such as extruded round (rod) or rectangular (bar) profiles of thickness 30-100 mm by lengths of up to several meters (Chang et al., 1986, 1987; Das, 1993). These profiles can be rolled to sheets of thickness 0.2-2.4 mm or forged into complex bulk shapes with and without SiCp reinforcements (Sec. 4.4.4; Das et al., 1989,1991,1992; Chang et al., 1988,1991). A large number of attractive improvements are offered by EA55RS, leading to the most superior Mg property profile yet reported. Typical strength values of mono-
4.4.2.2 State of the Art and Alloy Availability Figure 4-47 and Table 4-23 show a schematic diagram of the successive fabrication steps including the employed process details and the resultant properties of the RSP Mg-Al-Zn-base alloy EA55RS, respectively, made by the laboratory- and
Table 4-23. Engineering properties of the EA55A-RS alloy from Allied Signal in various forms (supplement to Fig. 4-47). No.
Condition
Corrosion in mm/year
UTS in MPa
TYS in MPa
in MPa
Elongation in %
1 2 3
As-extruded bar/rod As-extruded bar/rod T4-tempered bar low-high Ta
0.25 0.25 0.25
469-483 482-474 415-434
428-434 400-415 371-406
195 — —
10-14 12-14 14.7-24.5
4 5
As-rolled 2 h-temper at 325 °C
0.25 0.25
490-538 407
490-504 304
— —
4-6 14
a 30 min at 300 or 350 °C followed by water quenching; long transversal; T-L: short transversal.
b
in MPa m 1/2 6 6-8 L-T c :9 -15 T-L c : 6.5-7 7 —
rotating bending fatigue strength at 2800 Hz;
c
L-T:
166
4 Magnesium-Based Alloys
Principle Ribbon production
Process Details
1
Ribbon comminution
I Billet fabrication Mill processing
Planar flow casting: using inert gas shroud to stabilize fluid flow
Wheel speed: < 25 m/s Ribbon width: 25 mm Ribbon thickness: 20 - 30 Jim
Comminution to powder Sieving Matrix-base for composites
Sieve size: 250 - 500 jam (mean: 400 Jim)
I
Lab>: /
\
Pilot:
Vacuum degassing in can or vacuum hot pressing for 1 - 24 h at 200° - 300°C
Billet size 0 109 mm 1.8 kg
Billet size 0 279 mm 45 kg
1 - 4 h soaking & extrusion at 200° - 300°C, ratios 14:1 - 18:1
6501 Extrusion press Rod: 0 254 mm x Length 1.8 m
5500 t Extrusion press Rod: 0 75 mm x Length 6.7 m Bar: section 32 mm x 102 mm Sheet: thickness 0.4 - 2.4 mm
Sheet rolling at T = 250°C independent of 0 and speed of roll and reduction ratio Superplastic forging to complex shapes
Properties eg. EA55A-RS
Details of Material
Extrusion:
Corrosion Strength data Ductility Fracture toughness
Rolled sheet:
Corrosion Strength data Ductility Fracture toughness
7
I
Figure 4-47. Schematic diagram of the fabrication rate for RSP Mg-EA-base alloys and products using the Allied proprietary planar flow casting process. For 1-5 see Table 4-23.
lithic EA55RS extrusions at ambient temperatures are between 343 MPa (yield, TYS), 384 MPa (compressive, CYS), and 423 MPa (ultimate, UTS) at a 13% elongation-to-fracture EA, while the corrosion rate is about 0.25 mm per year corresponding to that of A12024-T6, but inferior to that of A17075. The special merits of EA55RS include superplasticity (at 300 °C: 436% EA) and a rotating bending fatigue strength of 180 MPa (for comparison QE22A: 100-105 MPa) (Das et al., 1992). The mechanical properties depend very much on the extrusion temperature and speed. The low-temperature (L-T) fracture
toughness Klc is about 6 MPa^/m and can be enhanced further, up to values of more than 15 MPa y/m at the expense of strength using a T4 temper between 300 and 400 °C (Das et al., 1992; Chang and Das, 1992 c). The creep behavior of this alloy is rather poor. However, EA55RS forgings with 10-30 vol. % SiC showed Young's modulus values of 71-92 GPa and thermal expansion coefficients of down to 12.8 x 10~ 6 /K. The galvanic corrosion behavior of this alloy has not yet been assessed. Current attempts to get RSP magnesium used in civil and military aircrafts and au-
4.4 Rapid Solidification Processing of Wrought Magnesium Alloys
tomobile applications involve the alloy EA55RS with the intention of encouraging the use of magnesium-base wrought alloys on a larger scale. Capacity and productivity of the process is not limited by PFC engineering and will respond according to market demand. The scope of the improvements and progress achieved by RSP of magnesium alloys in era 2 compared to conventionally processed magnesium alloys and the available aluminum alloys is as follows: l.An improvement of 40-60% in the room temperature ultimate tensile strength (UTS) over conventional ingot processed (I/M) magnesium alloys and the specific UTS of the strongest aluminum alloys (Fig. 4-48). 2. An increase in the ratio of the compressive-to-tensile yield strength (CYS/TYS) from 0.7 to a value of more than 1.1 (Fig. 4-49 and Table 4-24).
167
3. The resultant specific yield strengths ( T Y S / Q ) of RSP magnesium-base engineering alloys exceed those of I/M magnesium and aluminum alloys by 5298% in tension and 45-230% in compression. 4. The resultant elongations-to-fracture are within 5 and 15% for the as-extruded state and can be tailored to 22% by subsequent thermo-mechanical processing at the expense of moderate losses in strength. However, the corresponding strength values are still 150-200 MPa higher than for I/M magnesium alloys. 5. The atmospheric corrosion behavior of RSP magnesium alloys is in the range of the new, high purity, conventional alloys AZ91E and WE43, and the corrosion resistant aluminum alloy 2014-T6 (Fig. 4-50). The corresponding corrosion rates are two orders of magnitude smaller than for those magnesium alloys
RSP Mg alloys
300 ^5
Conventional & advanced Al alloys Conventional Mg alloys
PH
00
H
200
I H3 in
Figure 4-48. Specific ultimate tensile strength of chill-block quenched, RSP magnesium alloys developed from 1985 to 1992 compared with other light alloys. For details of alloy composition see Table 4-24.
168
4 Magnesium-Based Alloys
600
Commercial Mg Alloys _
500
Rapidly Solidified Mg Engineering Alloys
| U Yield Strength TYS |
Tensile Strength UTS
L J Elongation
400
Is
o
300
3 — 10
00 200 -
o 53
5
100
AZ91D-T6 ZK60A-T5
RS-AZ91+2Sr EA65-RS RS-AZ9 +0.9MM
L2
Figure 4-49. Mechanical properties of conventional I/M magnesium alloys and of RSP magnesium alloys with the latter trimmed to ductilities of more than 14%. For details of alloy composition see Table 4-24.
Conventional I/M Mg Alloys
0
o O
lO
1
1
1 B
3 M
< s r-|
11
ft <
3
I
• •
4 Li - 6.8 A 2.6
ft
2.3 Ca
1E-
ffi
too
EA65
"to
-
jiii
•::••
AZ9
o o
2
1I •i
5.8 O
a
CO
pm ||i
M2
"c5
DC c:
0
;-10.2A1-3.22
c 3 •22
-*
6Mn
— -
"E
Reference Al Alloys
-
g - 9 Al -
a> 4
Rapidly Solidified Mg Alloys
£ |
J<(J
^Hi
Jg «p M^
^ ^ ra ^
Figure 4-50. Corrosion rate of light alloys in 5% NaCl solution. Solid bars: novel magnesium alloy compositions afforded by rapid solidification. Unless otherwise indicated, data were obtained from weight loss measurements. 1} From electrochemical impedance spectroscopy.
4.4 Rapid Solidification Processing of Wrought Magnesium Alloys
that caused spectacular failures in the 1950s to the 1970s. 6. RS magnesium-base alloys show superior superplastic deformation behavior at temperatures above 100°C compared to other light alloys, and doubled resistance to fatigue as a result of substantial grain refinement compared to those obtained by ingot metallurgy. 7. The compatibility of RS magnesium alloys with reinforcements such as SiCp to serve as superior matrix material for resulting composites has been demonstrated (Das et al., 1992). Chill-block quenched, Mg-Al-Zn-base alloys are therefore primarily designed for applications at ambient temperatures and where compressive loading is of major concern, while galvanic problems may eventually be solved in other ways. It should be pointed out here that none of the research programs on RSP Mg-Li alloys have yet attained the specific strength levels of RS Mg-Al-Zn-base alloys. The specific strength is the particular property which formed the prime motivation for RS Mg-Li research projects. 4.4.3 Chill-Block Quenched MagnesiumBased Engineering Alloys: Extrusions 4.4.3.1 High Strength and High Corrosion Resistance at Ambient Temperatures
The evaluation of alloy EA55RS concluded a seven-year comparative alloy development program on RSP Mg-Al-Znbased alloys with five different quaternary and/or quinternary additions in order to optimize the corresponding multiphase intermetallic dispersions (Sec. 4.4.9.3). All of the alloys were hardened by grain refinement, by Zn in oc-magnesium solid solution, and by an Mg-Zn-based dispersion (Das et al., 1989). The choice and level of such additions were found, however, to be
169
crucial for the resultant resistance to microstructural degradation upon consolidation, so deciding the required extrusion speeds and temperatures as well as the increased resistance to corrosion, fatigue, and creep (Chang et al., 1986, 1987, 1988, 1991; Das, 1993; Das e t a l , 1989, 1991; Chang and Das, 1992 c). The alloys differed essentially between those containing Si and those containing rare earth elements (Table 4-24), resulting in additional dispersion hardening by either Mg2Si- and/or A12RE- (RE, e.g. Y, Nd) type dispersions. The former result in a corrosion rate of about 2.5 mm/year and the latter of 0.20.4 mm/year. Das et al. (1989) considered that the RE-containing dispersions or solid state precipitates improved the passivity of the surface film by reaction with atmospheric species, including halide anions from saline environments, but evidence is yet to be seen. The low difference in electronegativity between RE (and Y) and magnesium, however, will certainly support the inertness of the second phase particles, so reducing the tendency to pitting corrosion (Chang etal., 1986; Das etal. 1989). All of these RSP magnesium alloys had better specific strength values than the strongest aluminum alloys available today (Fig. 4-48), and their strength-ductility combinations were superior to the best 1/ M magnesium alloys (Fig. 4-49). Si-containing RSP Mg-Al-Zn alloys showed a tendency towards higher strength values at the expense of corrosion resistance (see above) and ductility (e.g. EA55A-RS with UTS = 516 MPa, TYS = 476 MPa, and elongation of 5%), as was the case when rare earth elements were added instead. Juarez-Islas (1991) reported a level of 3 % Si in a splat-cooled Mg-12A1- 2.5Zn alloy to allow for a maximum of some 30-50% increase in microhardness, compared to a
170
4 Magnesium-Based Alloys
Table 4-24. Chill-block quenched magnesium-base engineering alloys developed from 1985 to 1992a. Label
RS method
Allied Signal Al PFC A2 PFC A3 PFC A4 PFC A5 PVC
Designation
Composition in mass%, balance Mg Al
Zn
8.4 8.9 5.1 5.1 5.1
0.5 0.5 4.9 4.9 4.9
8.4 8.4 8.4 8.4 8.4 8.4 9 5 ZK60 + 3.2MM 5.1
0.7 0.7 0.7 0.7 0.6 0.7 3
EA55A-RS EA55B-RS EA65A-RS
Zr
Mn
Li
Y
Si
Ce
Nd
1.5 1.7 1.1
1.1
Pr
MM
Ca
Sr
5.3 5.0
6.7
Pechiney/Norsk Hydro PN1 PN2 PN3 PN4 PN5 PN6 PN7 PN8 PN9 Lookheed LI L2 L3 L4 L5
MS MS MS MS MS MS MS MS MS
RS AZ91 AZ91 + 1.5Si AZ91+0.9MM AZ91+2.0MM AZ91+2.3Ca AZ91+2.0Sr
0.13 0.13 0.13 0.13
1.5 0.9 2.0 2.3
0.13 0.6 0.6
2 6.5 2.5
0.2
3.0
3.2
Missile & Space Group (LMSG) MS MS MS GA GA
7 10 11 7.7 10.1
3 3 2 2.9 2.7
0.4 3
3
4 6 3
0.4 0.4
6 1.4
For comparison L4 and L5 are included here (see also Table 4-27).
splat-cooled EA65A-RS alloy that resulted in an exemplary combination of strength (UTS = 460 MPa, TYS = 432 MPa), ductility (20.6%), and corrosion resistance (12mpy). The best combination of strength, ductility, and castability was obtained for alloy EA55RS-B, which also showed a half an order of magnitude reduction in the fatigue crack propagation rate after heat treatment. EA55A-RS had the best damage tolerant properties includi n g ^ c values between 14 and 20 MPa y/m, improved notch impact strength ( + 25%), improved ductility ( + 33%), and superior strength (UTS: + 70 %, TYS: + 270%) compared to ingot processed AZ91HP (Chang et al, 1987). The fracture
surfaces of a corresponding fatigue specimen exhibited brittle cleavage features at low and intermediate AK values, and changed to ductile microvoid coalescence features at high AK values (Chang et al, . 1987). Dissolution, coagulation, and coarsening of the relatively coarse Mg 17 Al 12 equilibrium (Das and Chang, 1989; Chang et al, 1987) or chemically similar metastable phases (Li and Jones, 1992; Li et al, 1992 a; Ivanov et al, 1987; Graebner and Chen, 1987; Regazzoni et al, 1991; Sec. 4.4.9.3) resulting from a deliberate choice of the conditions of degassing, extrusion, and subsequent heat treatment, were considered (Chang et a l , 1987) to control the damage tolerant properties of
171
4.4 Rapid Solidification Processing of Wrought Magnesium Alloys
Density
UTS
TYS
CYS
in g/cm3
in MPa
in MPa
in MPa
Elongation in %
1.80 1.84 1.94 1.94 1.93
455 448 423 439 473
427 397 343 361 432
384 411 486
5.7 9.4 13 13.8 12.3
1.836 1.843
480-490 540 416 494 466/475 467 575 562
416 470 351 400 405/390 408 542-550 545
1.824 1.780
359 420
Young's modulus in MPa in MPa ^/nl in GPa
195
9-10 5.4 14.1 4.0 9.6 17 4.6-6.0 3.3
150
163 169
7.5-15 6.5 6.5
44.3 47.8
19
23 5.9
all as about for I/M AZ91
Corrosion rate in mm/year
2.4 0.38 0.36 0.28-0.38 0.30
1.3 1.9 8.4 7.8 0.20 0.58 0.15 0.58
444
1.73 1.99 1.89 1.98 1.91
383/380 468 380 410 380
346-309 431 310 345 315
337-381 488 397
this engineering RS magnesium alloy family (Sec. 4.4.9.1). Outstanding strength values were obtained by the European collaboration of Pechiney and Norsk Hydro by employing melt spinning or PFC (see below). The development concentrated on three modifications compared to that at Allied (Regazzoni etal., 1991; Gjestland etal., 1989, 1991; Nussbaum etal., 1990a, b, 1991; Project MAIE, 1990; Solberg et al., 1991; Lohne et al., 1991): (i) using AZ91 as the base and then adding Si and misch metal (MM); also Ca and Sr (Gjestland et al., 1989; Nussbaum et al., 1990b, 1991; Table 4-24), (ii) employing Mg-Al-Zn-based alloys as for Allied Signal, but with higher
4.9 14.9 4.3 5.5 4.0
47.5 42.7 46.2
0.27 0.47 0.61 1.3 1.5
aluminum levels ranging from 5 to 9%, reduced levels of Zn in the range of 0-3 %, and using Ca and Sr instead of an RE (Gjestland et al., 1989; Nussbaum et al., 1991); the Ca levels ranged from 1 to 6.5 % (Regazzoni et al., 1991), and finally, (iii) a ZK-type magnesium alloy (Mg-5Zn0.3Zr) with 3.2 % misch metal was evaluated by RSP (Nussbaum et al., 1991). The ribbons of thickness 30-100 |im and of width 12-15 mm (Gjestland et al., 1991) were directly cold compacted in either aluminum cans of diameter 45 mm or in press containers with a diameter of more than 120 mm to yield 70-90% dense billets prior to extrusion (Nussbaum et al., 1990a, b). Similar longitudinal and transversal
172
4 Magnesium-Based Alloys
strength values of RSP AZ91 flat products were considered to allow for abandoning the communition and degassing procedure prior to extrusion provided extrusion was done at temperatures between 250° and 350 °C and at a reduction ratio close to 20:1. An 800 t press was used for scalingup for direct extrusion with and without lubrication, or for indirect extrusion of billets of size 10 kg and diameter 150 mm. This scaling-up indicated the need for novel alloy compositions with temperature stable AlaXb dispersions (X = RE, Ca, Sr) in order to retain the laboratory-scale improvement in the scaled-up product form, such as rods of diameter 10-12 mm, bars of 48 mm x 15-25 mm, and I-beam and complex profiles with a satisfactory surface finish, the latter by using a multiple flow die (Nussbaum et al., 1990a, b). The refined multiphase microstructure did not show any degradation upon exposure for 24 h to 200 °C when Ca was absent, and for 24 h to 350 °C when Ca was added (Regazzoni et al., 1991). The employed ram speed was 0.5-3 mm/s and could be increased up to 5 mm/s when Ca was present. The application of RSP to the unmodified AZ91 alloy increased its TYS from 226 to 457 MPa ( + 106%) and its UTS from 313 to 517 MPa ( + 65%). The elongation-to-fracture of RS AZ91 ranged from 8.7 to 20.1 % depending on the extrusion temperature used (Regazzoni et al., 1991). As-extruded, 100 mm diameter rods with the novel RSP Mg-9Al-6.5Ca3Zn-0.6Mn alloy composition showed exceptional strength values of 575 MPa UTS; 542 MPa TYS, and 4.6 % elongation when made with direct lubrication, followed by RSP Mg-5Al-2.5Ca-3Sr alloy rods of the same diameter with 562 MPa UTS, 545 MPa TYS, and 3.3 % elongation (Nussbaum et a l , 1991). The novel RSP magnesium alloys by Pechiney/Norsk
showed corrosion rates of 0.2-0.6 mm/ year compared to 0.8 mm/year for RS AZ91, and were considered to have a comparable corrosion resistance to that of the aluminum alloy A380 (Fig. 4-50). The damage tolerant properties of RS AZ91 including the rotating bending fatigue strength
4.4 Rapid Solidification Processing of Wrought Magnesium Alloys
crb value was 163 MPa (Nussbaum et al., 1991). Lockheed's Missiles and Space Co. (LMSC) reported (Joshi and Lewis, 1988; Lewis and Joshi, 1989; Joshi et al., 1989) on an excellent combination of UTS = 468 MPa, TYS = 431 MPa, elongation = 14.9%, and CYS = 488 MPa for the as-extruded alloy L2 (Table 4-24). The most impressive feature of the RS-Mg alloys by LMSC was their ratios of compressive-totensile yield strength in the range of 1.051.13 (see Table 4-24). The CYS values of conventional I/M processed magnesium alloys do not usually exceed more than about 70% of the corresponding TYS due to excessive twinning upon compressive loading as a result of their large h.c.p. grains. The specific CYS obtained for alloy L2 was 450% larger than the specific CYS of AZ31B-F and 200% larger than that of ZK60A (see also Fig. 4-49), making these RS magnesium alloys highly attractive for space and aerospace applications where tensile and compressive properties at ambient or lower temperatures are critical and are of prime concern (Lewis and Joshi, 1989). LMSC observed emissions of CO and CO 2 above 250 °C and of hydrocarbons above 300 °C (Joshi et al., 1989), while Pechiney/Norsk had previously reported (Nussbaum et al., 1989 a) on the emission of water above 250 °C and of H 2 above 300 °C upon vacuum degassing of cold pressed billets prior to extrusion. However, Joshi et al. (1989) and Makar et al. (1988) found that unintentional oxidation and insufficient mechanical working of prior particulates gave rise to localized interparticle or grain boundary corrosion associated with particle loss, so exaggerating the weight loss (i.e. 8.9-11.4 mm/ year) in a gravimetric test, while corrosion rates obtained from electrochemical impedance studies were of the order of
173
those obtained by Allied Signal for EA55B-RS. LMSC was the first to consider using RSP magnesium alloys in aeronautical applications. Alloy LI (£ = 1.73 g/cm3) was considered (Lewis and Joshi, 1989) to meet all four failure categories required in order to replace the 7075-T76 aluminum alloy in a rocket-launched space support system, which is subjected to compressive and bending loading, and which involves sheet, thin plate, tubular, square- and tee-shaped extrusions. The reason for this was that alloy LI has an acceptable specific stiffness of 27.5 GPa cm3/g compared to 21.5 GPa cm3/g for alloy L2 and 24.4 GPa cm3/g for L3 (Lewis and Joshi, 1989). The relatively low specific stiffness values of the latter alloys were considered to be the major obstacle to replacing the aluminum alloys 6061-T6 and 8090-T8 in such space system applications, while aerospace applications would be more demanding (Lewis and Joshi, 1989; Thorbeck and Hehmann, 1996). 4.4.3.2 High Temperature Properties
Only very limited information is available in the literature on the high temperature properties of engineering RSP magnesium alloys. An RS-ZK-type magnesium alloy (Mg-5Zn-0.6Zr) with 3 % MM added showed superior resistance to secondary creep at 150°C compared to conventional AZ91 and ZE41 T5, for example (Table 4-25; Nussbaum et al., 1991). This improvement was attributed (Lohne et al., 1991) to the absence of any low melting point second phases on the grain boundaries of this ZK-type RS alloy. Temperature stable MgRE particles 0.1-0.3 |im in size were observed instead on the boundaries and in the interior of corresponding grains of size 0.6-0.7 jim. The creep stress
174
4 Magnesium-Based Alloys
Table 4.25. Secondary creep rates of engineering magnesium alloys at 150°C. Alloy
RS AZ91 I/M AZ91 RSAZ91+2%Ca RS Mg-5Al-5Ca-0.7Zn RSZK60 + 3%MM I/MZE41T5 b RSZK60 + 3%MM I/MQE22T6 b
Applied Second- Grain stress ary creep size in MPa in s" 1 in jim 50 50 50 50 50 100 120 125
5.1xlO~ 6 5.9xl0~ 8 l.lxlO~8 5.0xl0~ 9 1.2xlO~ 9
1.5 11 0.7a 0.7a 0.6a
3.5xlO~ 9 >60 2.3xlO~ 9 0.6a 9 50 0.6xl0~
a
Implies possible compositional effect upon secondary creep; b T5: 4h 215°C; T6: 16h 410°C/ quenched +4h215°C.
of EA55-RS after 100 h at 75 °C, however, decreased by 60-75% of its initial value, depending on the creep strain rate employed (Das et al., 1992). Nussbaum and co-workers (Gjestland et al., 1991; Lohne et al., 1991) reported on a more detailed analysis of the high temperature properties of RS AZ91. The tensile yield strength of RS AZ91, which had a grain size of 1.5 jim, decreased from 295 MPa at ambient temperatures to 150 MPa at 150°C (i.e. 50% less), while that of conventional AZ91 with a grain size of 12 jim decreased from 205 to 165 MPa (i.e. 20% less). The absence of grain growth in both conditions suggested that creep controls the high temperature properties of RS AZ91. The RS version showed a hundredfold higher secondary creep rate than conventional AZ91 under an applied stress of 50 MPa at 150°C. For a given grain size, the creep strength was observed to increase with decreasing aluminum content and the resulting decrease in the volume fraction of the relatively coarse (i.e. 0.5 |im) Mg 17 Al 12 grain boundary precipitates forming at about 150°C from supersatu-
rated solid solution (Hehmann, 1990 a). Obviously, such particles are too large to pin the motion of dislocations and grain boundaries, and so they increase the mobility of the grain boundaries instead due to their low melting point. The addition of 2.3% Ca to RS AZ91 confirmed the prime importance of chemistry (overall chemical composition) and the resulting diffusivity, growth and coarsening mechanisms in controlling the corresponding high temperature properties rather than grain size (alone): although RS AZ91+2.3% Ca had the smallest grain size of these three alloys (0.6 (im), its secondary creep rate was 400 times smaller than for RS AZ91 and 5 times smaller than for ingot processed AZ91 (Gjestland et al., 1991; Lohne et al., 1991). The corresponding tensile yield strength of 390 MPa at ambient temperatures decreased by only 65 MPa at 150°C. The stress relaxation after 100 h at 150°C followed the same order (Table 4-26). The fine trans- and intragranular dispersion of Al2Ca of mean size 0.05 jam was considered to pin dislocations and grain boundaries, suppressing both creep and grain boundary sliding even in the presence of temperature instable Mg 17 Al 12 phases (see Fig. 4-60). Relatively high initial stress relaxation rates in both RS versions indicated that grain boundary sliding controlled the beginning of high temperature deformation before transgranular creep became the rate-controlling mechanism. However, the rate-controlling creep mechanism of RS magnesium alloys remains to be identified in order to understand the effect of microstructural refinement and departure from equilibrium on the potential for the improvement of the high temperature properties and shapability close to product form of this novel category of wrought magnesium alloys.
4.4 Rapid Solidification Processing of Wrought Magnesium Alloys
175
Table 4-26. Comparison of high temperature properties of various AZ91 alloy versions at 150°C. Number
a
Alloy
Relative TYS-loss-compared to at ambient temperature
RS AZ91 I/M AZ91 RSAZ91+2.3%Ca
50% 20% 17%
Creep rate compared Stress relaxationa to alloy 3 400 times 5 times
88% 58% 45%
Ratio of initial stress (80 MPa) to remaining stress after 100 h at 150 °C.
4.4.4 Near-Net-Shape Forming of Chill-Block Quenched Magnesium Alloys The grain refinement obtained by RSP of magnesium alloys can lead to sufficient crystalline rotation upon plastic deformation at elevated temperatures to result in superplastic deformation, despite the limitations imposed by the h.c.p. crystallography (Chisholm and Busk, 1953; Chisholm and Hershey, 1956; Sec. 4.4.2.1). Excellent formability has been obtained with the RSP magnesium alloys by Allied and by the Norsk-Pechiney-collaboration due to the combination of low flow stress and high ductility at temperatures as low as 150 °C, making such alloys useful for nearnet-shape operations including rolling and forging (Das et al., 1990, 1991; Solberg etal., 1991; Raybould et al., 1991; Chang and Das, 1992 a, b). 4.4.4.1 Forging of Monolithic EA55RS, the Resultant SiCp Composites, and RS AZ91 Das et al. (1990) reported superplastic forming and forging operations using vacuum hot-pressed billets, directly canned containers, and extrusions of alloys EA55B-RS, EA65A-RS, and R S - M g 5Zn-5Al-2.7Ce at temperatures ranging from 150° to 275 °C in order to arrive at complex component shapes of minimum thickness 1 mm. This is an up to 150°C
lower forming or forging temperature than is required for the I/M ZK60 alloy, for example. High precision components without cracks were obtained by closed-die forging at temperatures as low as 160°C when the ram speed was of the order of 0.01 mm/s. Higher forging temperatures of 220-240 °C allowed ram speeds of up to 0.21 mm/s. Exceptional UTS values of 510-540 MPa, TYS values ranging from 450 to 484 MPa, and elongations from 6.8 to 9.1% were obtained after forging EA55B-RS extrusions made at low extrusion speeds, while prior extrusion at higher speeds was found to lead to some softening" upon subsequent forging, together with ductilities of 10-13% (Das et al., 1990). An alternative route was the multiple forging operation using closed- and opendie steps for hot isostatically pressed or cold-pressed and then sintered EA55B-RS billets in order to produce pancakes of diameter 14 cm x height 2 cm at forging temperatures between 200 and 300 °C (Raybould et al., 1991). A number of closed-die forging steps at thickness reductions of about 20-25% were followed by a final open-die forging step with a reduction in thickness of 50%. This procedure did not result in any serious cracking. Strength and ductility values of EA55B-RS were 438-504 MPa UTS, 400-469 MPa TYS, and 1.4-6.3% elongation-to-fracture, respectively, on using a sequence of five forg-
176
4 Magnesium-Based Alloys
ing steps. The ductility was substantially improved at the expense of minor strength losses by additional working via one more forging step in the same temperature regime, or by reducing the number of forging steps but increasing the forging temperature at 300°C (Raybould et al., 1991). The microstructure of such monolithic forgings consisted of grains of size 0.21.0 jam, and of intermetallic phases of size between 0.03 and 0.07 jim (Raybould et al., 1991). Hot pressing at temperatures between 250° and 500 °C was reported (Das et al., 1992) to result in sufficient bonding strength between comminuted EA55RS matrix powders and SiCp reinforcing particulates without a detrimental coarsening effect on the intermetallic matrix dispersion. Subsequent extrusion and annealing of EA55RS with 10 vol.% SiCp for 0.5 h at temperatures between 350 and 500 °C resulted in a UTS value of 570 MPa, a coefficient of thermal expansion a of 19 x 10~6/K, and a Young's modulus of 72 GPa. Forgings of EA55RS with 30 vol.% SiCp show a values of 12.8 x 10" 6 /K and a modulus of 85 GPa (see also Sec. 4.5). Nussbaum et al. (1990 a) reported critical strain values for the onset of vertical surface tearing and cracking of the RS AZ91 alloy extruded at 250 °C. Note that surface tearing was one of the reasons why the pioneering work by Dow Chemical on RS magnesium engineering alloys ceased in the early 1960s (Couling, 1985). The observed critical values ranged from a strain of about 60 % at forging temperatures of 250°C and a strain rate of 1.5 x 10" 3 s" 1 down to about 20 % at a forging temperature of 150°C and a strain rate of l . S x l O ^ s " 1 under a given set of conditions. Friction between metal and die and the resultant magnitude and distribution of stress in the metal was of great importance, as it was upon rolling between metal
and rolls (Sec. 4.4.4.2). At 250 °C, for example, lubrication of RS AZ91 allowed for an up-setting of 215% without sacrificing any of the hardness increment obtained by RSP (Nussbaum et al., 1990b). 4.4.4.2 Sheet Rolling of EA55B-RS Sheet rolling is a very attractive alloy conversion route. It is per se a high productivity process, since it drives relatively large preforms of metal under high-speed and continuous conditions directly into the final product shape. However, conventional magnesium alloys possess a quite narrow temperature interval between the minimum temperature to avoid cracking upon rolling and the maximum temperature to avoid alloy softening (Emley, 1966), and thus the corresponding market has almost entirely disappeared (see Sec. 4.1). Rolling of EA55RS-B is performed using a mill with two vertically superimposed steel rolls of 13 cm diameter at 200-300 °C applying preferably 5-20 passes with a thickness reduction of 4-10% per pass (Chang and Das, 1992 a). Cracking upon rolling at higher ram speeds of EA55RS-B was minimized by using a rolling temperature of 250°C (Chang and Das, 1992b). Unlike for commercially available, conventionally processed magnesium alloys, EA55RS-B can then be hot rolled down to a thickness of 0.2 mm without cracking (Chang et al., 1991; Das et al., 1990; Chang and Das, 1992 a). EA55RS-B sheet was found to develop a strong basal {0001} texture in parallel to, or near the rolling plane, which was nine times stronger than for the extruded preform, as well as an unusual {1013} twinning texture three times stronger than in the extruded preform (Chang and Das, 1992 a, b). The texture consisted of intragranular subgrains of size 0.1-0.2 |im and dispersoids of size
4.4 Rapid Solidification Processing of Wrought Magnesium Alloys
0.02-0.04 |im, and a dislocation network that was not present in the employed extruded preform. This texture results in high strength, but low ductility. The ductility was improved, however, by post-fabrication heat treatment at 325-350°C from about 6-14%, but it was not improved by increasing the rolling temperature (Table 4-23, Nos. 4 and 5). Annealed EA55RS sheet with UTS = 407 MPa and TYS = 304 MPa is about + 4 0 % stronger than the strongest conventional magnesium alloy sheet made from the I/M AZ31B-H24 alloy (ductility = 15%). These results were attributed (Chang and Das, 1992 a) to dynamic recovery which underlies, on the one hand, the embrittlement upon rolling. On the other hand, dynamic recovery also isolates crack cavities that may form upon rolling, thus avoiding catastrophic failure as occurs in conventional alloys under corresponding conditions, and it allows for softening, so eventually tailoring the final product (Table 4-23). 4.4.4.3 Superplasticity The microstructure of RS Mg-Al-Znbase alloys allows for significantly higher superplastic forming rates than for most of the other light alloys. At 150°C and a strain rate of more than l x l O ~ 3 s ~ 1 , for example, as-extruded EA55B-RS and EA65A-RS alloys showed elongation-tofracture values between 190 and 220% (Das et al., 1990) allowing for forging operations of highly complex components to be carried out without the occurrence of cracking (Sec. 4.4.4.1). This increase in ductility at 150°C was associated with a reduction in the strength to a third or less of the corresponding room temperature strength. The strain rate sensitivity of EA55B-RS increased dramatically at temperatures of more than 100 °C (Das et al.,
177
1990). Fabrication of complex shapes was also reported (Chang and Das, 1992 b) by superplastic forming of EA55B-RS sheet (Sec. 4.4.4.2). The corresponding values of elongation-to-fracture ranged from 376 to 436% at temperatures between 275 and 300 °C using a strain rate of not less than 0.1 s" 1 (Chang and Das, 1992a, b). These values were obtained for material made at the pilot production plant (Fig. 4-47). The optimum temperature for superplastic forming of EA55B-RS sheet was found to be 300 °C. At a forming temperature of 275°C and a strain rate of 0.01 s" 1 , the elongation was still around 300 % for samples made from a laboratory scale batch, while it was 250 % from pilot production scale batches (Chang and Das, 1992 a). Some strain hardening as well as yield point effects was observed at strain rates of more than 0.01 s" 1 , indicating dislocation pinning by solute atoms. However, no grain coarsening was observed during such operations. RS AZ91 showed elongations of more than 1000% at temperatures between 275 and 300 °C compared to 240 % for the corresponding ingot material (Lohne et al., 1991). The observed coarsening of the original microstructure of RS AZ91 (i.e. grains of uniform size 1.2 + 0.4 \xm decorated by Mg 17 Al 12 particles of size 0 . 1 0.3 jim; Sec. 4.4.9.3) upon exposure for as long as 21 h to 300 °C was limited to a grain size of 1.9 + 0.7 jim without affecting the superplasticity. The maximum strain rate sensitivity for both alloys was about 0.6, but for I/M AZ91 it was observed at a two orders of magnitude lower strain rate than for RS AZ91. Increasing porosity around temperature instable, Mg 17 Al 12 grain boundary particles with increasing strain rate employed, and an 18-36% lower activation energy for superplastic forming compared to that for the self-diffusion
178
4 Magnesium-Based Alloys
of magnesium were considered to show that grain boundary diffusion was the controlling deformation mechanism for RS AZ91 (Sec. 4.4.3.2). This porosity led to a degradation in the strength to 385 MPa and the elongation-to-fracture to 17% at ambient (Solberg et al., 1991). The development of more temperature stable, wrought magnesium alloys by RS methods does therefore not necessarily impose limitations on the corresponding superplasticity and mechanical properties at lower temperatures. In a different study on a conventionally processed, high temperature magnesium alloy, Zelin et al. (1991) showed that grain boundary sliding (GBS) controlled superplastic deformation under the optimum conditions of deformation. The reduction in grain size from 100 to 10 pm of high temperature Mg-l.5Mn-0.3Ce alloy sheet obtained by cold rolling and annealing was reported to halve the corresponding flow stress at a temperature of 400 °C and a strain rate of4xl0~~ 4 s~ 1 , and to reduce the density of dislocations by a factor of 16. The optimum ratio of the effective stress required for GBS to the inactive transgranular stress consumed for dislocation movement in Mg-l.5Mn-0.3Ce alloy sheet was found to be at around 400 °C, and was considered to increase from a ratio of 50% to 50% for the 100 Jim grained version to 70% to 30% for the 10 Jim grained version. At 350 °C, the optimum ratio was assumed to account for 65 % to 35 %. It is possible that such optimum conditions for superplasticity would cause softening of EA55RS, thus losing substantial strength increments obtained by the PFC process, since the microstructure of EA55RS consisted of temperature instable, icosahedral Mg 32 (Al,Zn) 49 -type grain boundary phases which dissolve at 300 °C similar to Mg 17 Al 12 in RS AZ91 (see Sec. 4.4.9.3).
4.4,5 Gas-Atomized and Spray-Formed Magnesium-Based Engineering Alloys
Promising modifications of processing include spray deposition techniques such as the Osprey process and liquid dynamic compaction (LDC), in which the fragmented volume is exposed to an inert gas atmosphere for less than a fraction of a second prior to consolidation of the corresponding particles on a partially molten, preform surface. The fragmentation in situ consolidation procedure was sufficient for the magnesium alloys to retain refined grains of the order of 1-50 pm in size depending on the alloy constitution and spray-forming conditions used (Faure et al., 1991; Lavernia et al., 1987; Vervoort and Duszczyk, 1991; Elias et al., 1992). Spray deposition circumvents a number of processing steps required for manufacturing RSP magnesium alloys from chillblock, melt-spun (see above), and gasatomized (see below) materials, such as comminution and/or canning, vacuum degassing, etc. Spray-formed (SF) materials usually contain the lowest levels of RS process contaminants such as oxygen, resultant oxide dispersions, hydrogen, etc., so allowing for a major improvement in the fracture toughness K1C along with sufficient increments in the other mechanical (strength, ductility) and electrochemical properties compared to the conventional ingot route. Faure et al. (1991) reported fracture toughness values of 30 and 35 MPa ^/m for spray-formed Mg-7Al-4.5Ca-l.5Zn1.0RE and Mg-8.5Al-2Ca-0.6Zn-0.2Mn, while the UTS (TYS) was 480 (435) and 365 (305) MPa and the elongation-to-fracture was 5% and 9.5%, respectively, for these alloys. Both the fracture toughness and the resistance to 107 rotary bending cycles were superior to those for I/M AZ80
4.4 Rapid Solidification Processing of Wrought Magnesium Alloys
and RS AZ91 + 2Ca made from melt-spun ribbon. The corresponding microstructure consisted of grains of size 3-25 jim including preferential grain boundary precipitates corresponding to Mg 1 7 Al 1 2 , Al 2 Ca, MgRE, and A1RE (Faure etal., 1991). Much better ductility without sacrificing strength was obtained (Lavernia et al., 1987) for LDC Mg-5.6Zn-0.3Zr and LDC Mg-8.4Al-0.2Zr compared to ingot processed ZK60 and AZ80 in the as-extruded and aged conditions (Table 4-25). The combination of strength and ductility for LDC Mg-5.6Zn-0.3Zr was as for meltspun ZK60 (Anonymous, 1987) despite its lower Zr content and less extreme conditions of solidification. LDC Mg-Zn-Zr showed recrystallization to result in grains of size 5 jim on extrusion, followed by coarsening to 30-40 jim after solutionizing for 1 h at 500 °C and ageing for 48 h at 130°C. However, no recrystallization and coarsening was found for the LDC M g Al-Zr alloy even after 5 h solutionizing at 413°C and 20 h ageing at 205 °C, due to more stable precipitates such as Al 3 Zr which were not present in the former alloy. Improvements in strength ( + 40%), ductility (from 3 to 10%), and resistance to corrosion by a factor of three were also reported (Vervoort and Duszczyk, 1991) for the spray-formed magnesium alloy QE22 compared to the corresponding I/M alloy. Elias et al. (1992) observed solid state precipitates as fine as 30 nm after co-spraying Mg-Mn alloys with 56 \imsieved, pure aluminum and aluminum alloy powders, while spray-formed WE54 showed equi-axed grains of the order of 10 jim. The porosity of all of these alloys did not exceed 5 % in the as-sprayed condition. Improved atomizer design, resulting in finer and more spherical powders with a narrower particle size distribution (Hop-
179
kins, 1991-1993), and the advent of spray deposition techniques have nonetheless stimulated interest in the development of both monolithic and composite magnesium-base alloys by gas-atomization (GA) methods, despite their less stringent conditions of heat extraction than afforded by chill-block methods. It has been proposed (Unal, 1992) that there is no basic difference in the median diameter of magnesium-base and aluminum-base powders provided they are made under identical conditions. Inert gas atomization of alloys L4 and L5 resulted in substantially reduced corrosion rates (1.3-1.5 mm/year) over corresponding chill-block-made versions without localized corrosion at equal (L5 versus L3) or some 10% lower (L4 versus L2) strength values (Sec. 4.4.3.1; for alloy designations see Table 4-24). The effect of the interparticle oxide distribution upon the corrosion behavior of RS-Mg alloys was therefore considered (Joshi et al., 1989,1992) to be more important than the effect of precipitation, though such improvements are not inherent to the type of RS processing employed, as was indicated by the results with melt-spun RS magnesium alloys by Allied and Pechiney/ Norsk (see above). Though spherical powders may allow a better redistribution of the surface oxides resulting from particulate surfaces, degassing methods (in a can or not) and the degassing temperature of cold-pressed RS magnesium alloys remained the crucial factors to be optimized for the best corrosion behavior of this family of chill-block, melt-spun magnesium engineering alloys (Joshi etal., 1992; Baliga etal., 1992; Sec. 4.4.3.1). Interesting results were obtained by ultrasonic GA of the Mg-3.2Nd-l.lPrl,5Mn alloy and subsequent consolidation at low temperatures. Krishnamurthy et al. (1988, 1989) reported substantially better
180
4 Magnesium-Based Alloys
Table 4-27. Gas-atomized (GA) and Spray-formed (LDC & SF) magnesium-base engineering alloysa. UTS in MPa
TYS in MPa
Elongation in %
LDCMg-8.4Al-0.2Zrb LDCMg-5.6Zn-0.3Zrc SF Mg-7Al-l.5Zn-4.5Ca-l.0RE SF Mg-8.5Al-0.6Zn-2Ca-0.2Mn SF QE22
351 354 480 365 350
253 303 435 305 290
18 14 5 9.5 10
Ga Mg-3.2Nd-l.lPr-l.5Mn Ga AZ91 Ga ZE63 Ga ZK60 T6 GA QE22 GA Mg-7.9Al-0.7Si
427 400 430 403 415 405
420 350 410 376 380 291
5.1 10 4 17 4 19
Method and alloy
a
In the as-extruded state, if not indicated otherwise; atl30°C.
tensile strength at room temperature and improved corrosion resistance compared to the strongest I/M alloy ZK60 (Table 4-27), although this alloy did not contain any of the classical solid solution strengthening elements such as Al and Zn. Satisfying interparticle bonding without porosity was attained after degassing for 1 h at 250 °C, preheating for 2 h at 250 °C, hot isostatic pressing for 6 h at 250 °C, followed by preheating for 2 h at 250 °C and extrusion at reduction ratios between 12:1 and 20:1 at 100-250°C or 8:1 at 150250 °C. Outside these conditions, however, no satisfying consolidation was possible for GA Mg-3.2Nd-l.lPr-l.5Mn. Kainer (1989, 1990, 1991a) reported strength improvements of 40-300% for as-extruded GA over the I/M magnesium alloys AZ91, ZE63, and QE22 (Table 4-27). GA AZ91 showed (Kainer et al., 1991) a major improvement in ageing response after 3 h exposure to 175 °C, resulting in a maximum UTS value of 400 MPa and a maximum TYS value of 350 MPa in the as-extruded condition. No studies on
K1C in MPa y/ni
Corrosion rate in mm/year
30 35
0.46 0.15 3.3
after temper for 20 h at 205 °C;
0.28
14.7 c
after temper for 48 h
grain size were undertaken to interpret these results on a more quantitative basis (cf. Sec. 4.4.9.2). The ductility of these GA magnesium alloys either decreased or was only slightly improved compared to corresponding I/M versions. A refined dispersion of intermetallics and oxides (the latter not present in cast ingot materials) was suggested (Kainer, 1991a) as suppressing recrystallization of GA QE22 upon extrusion of 350 °C. Subsequent thermo-mechanical treatments improved the mechanical properties of these PM magnesium alloys in monolithic form (Kainer, 1989, 1990). Reinforcement with SiC particles resulted in an increase of 1.4 GPa in the elastic modulus per vol.% SiC for GA AZ91, and of 0.7 GPa per vol.% SiC for GA QE22 (Kainer et al., 1991; see also Sec. 4.5 and Table 4-29). Finally, gas atomization of Mg-8Al-0.8Si was reported to result in a TYS = 291 MPa (grain size 2 jim), UTS = 405 MPa, and 19% elongation, but the corresponding corrosion rates were rather poor (Garboggini and McShane, 1992).
181
4.4 Rapid Solidification Processing of Wrought Magnesium Alloys
4.4.6 Rapid Solidification Processed (RSP) Magnesium-Lithium Alloys 4.4.6.1 State of Development The need to overcome the low strengthening response of Li in h.c.p. oc-Mg, to reduce alloy softening at higher Li levels for (oc+p) and b.c.c. (3-Mg alloys, to improve the resistance to overageing, creep, stress corrosion cracking, atmospheric corrosion, and intragranular embrittlement upon conventional processing along with precipitation hardening additions such as Al, Zn, and Ag, and to replace toxic additions such as Th and Cd has stimulated interest in RSP as a method of producing Mg-Li based alloys. The prime objective of RSP was to introduce finer and thermally more stable intermetallic dispersions than are possible by conventional means. Despite their low densities, however, none of the current RSP Mg-Li alloys can compete with the new, highstrength RSP magnesium alloys (see Sec. 4.4.3, and Figs. 4-48 and 4-49). This is particularly evident for their resistance to plastic deformation (compare the TYS values in Tables 4-24 and 4-28). Meschter and O'Neal (1984), and Meschter (1986) reported a 50-60% increase in strength of the twin-roller quenched (a + P) Mg-9Li-2Si (or 2Ce) alloy over conventionally cast Mg-9Li (Table 4-28). Half of this increment was obtained by rapid solidification alone, the other half resulted from a finer dispersion of Mg2Si and Mg 9 Ce. The corresponding Hall-Petch grain boundary strengthening coefficient ky is, however, very low (see Sec. 4.4.9.2 and Fig. 4-59). The improvement in strength was even more evident at 150 °C where the increment was 3 times as much for RS Mg-9Li with and without 1 % Si and 4 times as much with 1.7 % Ce. The fine, intermetallic dispersions trigger
Table 4-28. Rapidly solidified magnesium-lithium base alloy extrusions. Method a and alloy
UTS TYS Elongation in MPa in MPa in %
I/M Mg-9Li TR Mg-9Li TR Mg-9Li-2Si TR Mg-9Li-2Ce
124 155 196 188
103 131 160 157
14 12 17 18
I/MMg-10Li-5Sic MSMg-10Li-5Sic I/MMg-8Li-5Al c MSMg-8Li-5Alc
128 205 168 208
98 200 145 200
22 0.6 b 8 lb
a TR = twin-roller quenched, MS = melt spun; b inhomogeneities in the specimens; c data from Schemme (1993 b).
recrystallization so that further grain refinement occurs upon consolidation, but they were observed to suppress dynamical recrystallization upon uniaxial loading at higher temperatures due to stabilizing the micro structure for exposures up to 300 °C. Grensing and Fraser (1987) attributed a decreasing lattice parameter for ceritrifugally gas-atomized Mg-8Li-1.5Si to supersaturation of oc-Mg with Li, although Si might have coexisted with Li in the extended solid solution, as was supported by the presence of very fine silicides after overageing for 2 h at 300 °C. Kalimullin et al. (1986, 1988) and Kalimullin and Berdnikov (1986) reported using laser surface treatment of the Mg-8Li-5Al-4CdlZn-0.4Mn alloy (Soviet designation MA21) to improve its resistance to creep and to corrosion (one order of magnitude in 3 % NaCl solution) due to a fine quasi eutectic surface structure. The microhardness of the laser-treated, binary Mg-8Li alloy was found to increase by 40% over the corresponding underlayer, and by more than 600% if prior-clad with pure aluminum (Schemme, 1993 a).
182
4 Magnesium-Based Alloys
Table 4-29. Mechanical properties of magnesium metal matrix composites (MMCs). MMC
Modulus in GPa
P55/AZ91C (40 vol. % unidirectional) FT700/AZ61 (57 vol.% unidirectional) Mg/T300 Mg/T300 Mg-4Al/T300 Mg-4Al/T300 (all 30 vol.% unidirectional) Mg/P55 ZE41A/P55 AZ91C/P55 Mg/P55 ZE41A/P55 AZ91C/P55 (all 40 vol.% unidirectional) AZ31B/SiCp 20 vol.% Mg-Zn6/SiCp 20 vol.%
172
Temperature in MPa in MPa in °C
UTS
829
827
RT
968
RT
325 522 328 645
RT RT RT RT
659 279 587 12 19 45
RT RT RT RT RT RT
longitudinal longitudinal longitudinal transversal transversal transversal
Chin and Nunes, 1988
341 215 466 390 281 173 398 346 159 217 134 258 152 280 253 280 252 330 190
RT 150 RT 95 150 200 RT 100 200 RT 200 RT 200 RT 200 RT 200 RT 200
stirring + extruded
Mikucki et al., 1990 Mikucki et al., 1990
159 204 184 20 25 28 79 68 75 66 59 50
ZC71/SiCp 12 vol.%
Mg/Al 2 O 3
50
Mg-2.5Ag/Al2O3
53
Mg-2.5Nd/Al2O3
60
QE22/A12O3 all 20 vol.% Saffil QE22/(A12O3 10 vol.% + SiC 15 vol.%)
58 77
Comments
YS
270 167 418 332 227 133 370 230 121
290
Schemme and Hornbogen (1990) and Schemme (1993 a) reported a maximum in the microhardness at 5% Li for binary Mg-Li alloys in both the melt-spun and the ingot processed state, i.e. at the borderline of the oc to the (oc + (3) two-phase field. The maximum increment amounted to 24% for the RS over the corresponding I/M Mg-5Li alloy. Alloy softening at
References
Goddard et al, 1986 Rabinovitch etal, 1992 fibers surface treated Hall, 1991 no surface treatment fibers surface treated no surface treatment
stirring + extruded
stirring + extruded
Wilks and King, 1992
squeeze casting
Kainer, 1991b
squeeze casting
Schroder and Kainer, 1991
higher levels of Li was particularly evident in the (oc + (3) two-phase field for both RS and I/M processing, while alloy softening inside the b.c.c. |3-Mg concentration regime was relatively small (Schemme and Hornbogen, 1990; Schemme, 1993 a). These results confirm that the b.c.c. crystal structure is responsible for the large amount of alloy softening at higher Li con-
4.4 Rapid Solidification Processing of Wrought Magnesium Alloys
centrations, while the Li concentration itself may play a minor role above 11 % Li or so. Binary RS Mg-Li alloys with more than 11 % Li did not show more than a 10% hardness increment over corresponding I/M alloys, making it questionable whether b.c.c. p-Mg alloys are suitable candidates for demanding structural applications (see below). Schemme and Hornbogen (1990, 1992) and Schemme (1993 a, b) reported a 60% strength improvement for as-extruded RS Mg-10Li-5Si ribbon over the corresponding I/M alloy and a 24% increment resulting in 210 MPa for RS Mg-8Li-5Al (Table 4-28). It was interesting to note that an addition of 5 % Si and Al shifts the maximum strengthening response from the a-phase field into the (a + (3) two-phase field at 10 and 8% Li, respectively (Schemme and Hornbogen, 1992; Schemme 1993 a), while the strength increments were observed to decrease for lower and higher levels of Li in both ternary systems. The shift of the onset of alloy softening to higher Li levels, i.e. the (a + p) two-phase field, was considered to result from supersaturation of b.c.c. p-Mg with Si, resulting in Mg2Si particles 2030 nm in size (Schemme and Hornbogen,
500
100 50
M l
10
1
100 -
1992). A more thorough evaluation suggests, however, that grain size is an equipotent strengthening mechanism in the Mg-Li system (Fig. 4-51, see also Sec. 4.4.9.2). 4.4.6.2 Magnesium-Lithium-Hydrogen Alloys Haferkamp etal. (1990, 1992, 1993) have recently announced a large potential for the development of b.c.c. P-MgLi alloys by gas atomization using hydrogenenriched Mg-16%Li (40at.%) alloy melts. A concentration of 0.12% (2 at.%) hydrogen was reported (Haferkamp et al., 1990) to form a dispersion of Li hydrides and/or hydrided Li zones of the order of 10-40 jim in size in chill-cast Mg-16Li. While chill-casting of the Mg-6Li-0.12H alloy did not improve its strength, gas atomization of Mg-16Li-0.12H resulted in a minor increment of 20 MPa compared to the 100 MPa of Mg-16Li. Haferkamp etal. (1990, 1992, 1993) considered (i) a high mobility of hydrogen atoms in magnesium and (ii) the affinity of hydrogen to lithium as instrumental in trapping and tying-in the likewise very mobile Li atoms.
5
'
3
'
'i
90 80
^
^
5Li
70 . -"^
60
3Li
50 40
-
20Li i
i
i
0.1
0.2
binary hcp-oc-MgLi# binary bcc-p-MgLi © LA141 # EZ32 © i
0.3
I
0.4
183
I
0.5
Figure 4-51. Hall-Petch-type relationship of magnesium-lithium alloys for the Vickers hardness number (data obtained using a 25 g load) as a function of the reciprocal value of the square route of the grain size d~l/2. For a given structure of the binary alloys this relationship appears to be independent of the lithium concentration, as indicated in mass%. For comparison, data for as-splatted EZ32 are included (Hehmann, 1984; Schemme, 1993 a).
184
4 Magnesium-Based Alloys
Based on this hypothesis first forwarded by Borbe and Erdmann-Jesnitzer (1983), Haferkamp postulated that bulk (volume) diffusion assisted Li fluctuations to control the room temperature properties of b.c.c. p-magnesium alloys, and that an increased addition of hydrogen upon gas atomization renders items (i) and (ii) instrumental for more effective hydride-dispersions, so allowing for strength levels of 300 MPa (Haferkamp et al., 1990). However, the maximum strength of a gas-atomized and as-extruded Mg-16Li-4.6Al alloy was only 200 MPa (elongation 20%) compared to 264 MPa (9 % elongation) for the corresponding I/M alloy, which was considered as a base system for higher order alloys using hydrogen and PM technology (Haferkamp et al., 1992). Portevin-LeChatelier-type plastic flow has been observed in many other alloy systems without attributing the rate-controlling mechanism to diffusion. Haferkamp's postulations leave a number of concerns: 1. All Mg-Li alloys including the P-b.c.c. variant were reported to show positive Hall-Petch proportionality constants ky and intercepts (7^ cpMgLi X J ^ 8 and ^b.c.c.MgLi > ^.Mg indicating that the corresponding mechanical properties are controlled by plastic flow via slip and not by diffusive flow via climb, which would require the ky value to be negative (Hehmann, 1995 b; and Sec. 4.4.9.2). 2. The b.c.c. P-Mg-16Li-0.12H alloy shows the lowest of all of the reported ky values (Fig. 4-59), indicating a decrease rather than an increase in the boundary strengthening response due to the introduction of hydrogen in Mg-Li alloys when compared with Mg-Li alloys without deliberately added hydrogen (Sec. 4.4.9.2). The ky value of b.c.c. pMgLi alloys is in the range of that of
pure Al, but nobody has yet attributed the low ky value of pure aluminum to diffusion. 3. A strength of around 300 MPa would require a grain size in the range of 0 . 1 0.3 jim for the ky values obtained for b.c.c. P-MgLi and (oc + p) alloys (Fig. 4-59). The smallest grain size yet reported for a b.c.c. P-MgLi alloy was of the order of 10 \xm using chill-block methods (Das et al., 1988; Schemme and Hornbogen, 1992), which provide far more extreme conditions than gas atomization. 4. Diffusion may be the rate-controlling factor for the coarsening of second phase dispersions, but the mechanical properties are usually controlled by the way the matrix dislocations respond to them and also to other obstacles, including the crystal structure itself. Relevant hydrides, however, have yet to compete more effectively with those intermetallics resulting from silicon and rare earth metal additions to magnesium. 5. Diffusion, however, is related to the atomistic properties and it is therefore proportional to the mean melting point of the involved constituents. Phase diagram considerations (see Haferkamp et al., 1990, 1992, 1993; Borbe and Erdmann-Jesnitzer, 1983) do in fact obscure atomistic properties such as diffusion. An increased use of hydrogen could thus easily increase and not decreae the susceptibility to thermal activation of the rate-controlling deformation mechanism (Schemme and Hornbogen, 1990, 1992; Schemme 1993 a; Haferkamp et al., 1990, 1992, 1993; Borbe and Erdmann-Jesnitzer, 1983). 6. Body-centered cubic P-Li is thermodynamically speaking a congruently melting b.c.c. compound with a large solubility range, and is not diffusive matter,
4.4 Rapid Solidification Processing of Wrought Magnesium Alloys
though the degree of ordering has not yet been identified. The resultant increase in the modulus of elasticity over h.c.p. oc-Mg should affect the stacking fault energy and/or the preferred system of slip, but this crystallography should not trigger per se nonconservative diffusive flow by climb. Unless thermally activated variants of slip (cross slip, dislocation kinks), or ordering phenomena, or any interrelation between them can be excluded, current research into the use of hydrogen for more advanced b.c.c. (3-MgLi alloys remains pure speculation. In h.c.p. a-Mg and corresponding h.c.p. oc-MgLi alloys, however, diffusion and climb phenomena do certainly not control the mechanical properties (Hehmann, 1995 b).
melt-spun Mg 70 Zn 30 ribbons. This was twice as strong as today's strongest commercial magnesium alloy (Sees. 4.3.1.2 and 4.3.2). Together with the attractive corrosion features of hypersaturated solid solutions first found by Hehmann et al. (1989), metastable magnesium alloys form a novel route for the development of superior light alloys. The major drawbacks yet encountered with amorphous and metastable crystalline or quasi-crystalline magnesium alloys include susceptibility to natural ageing (Fig. 4-52; Hehmann, 1990 a) and embrittlement (if the weight penalties are acceptable), their generally low thermal stability, and large weight penalties in par-
Evidently, RSP of magnesium alloys has not yet been explored to its full potential, but it has opened the door to a more effective hierarchy of future tasks. However, unless the low strengthening response and the rate-controlling deformation mechanism in b.c.c. (3-MgLi are understood, it seems questionable whether the Mg-Li system and in particular the b.c.c. variant can render magnesium capable of replacing plastics.
400
300
500 Temperature in K
.312
4.4.7 Amorphous Magnesium Alloys
Ultra-high strength amorphous Mg structures and ultra-high resistance to corrosion of hypersaturated a-Mg solid solutions (Sec. 4.4.8) are increasingly attracting attention to nonequilibrium processing for the development of superior magnesium alloys in the metastable state. The great capacity of magnesium to form metastable alloys without crystalline longrange order had been demonstrated by Calka et al. (1977) who reported exceptional UTS values of up to 840 MPa for
185
168
•E - 1 0 CD
I Q. g>
\ -
-20 -
1 LU
If
1,
96
ill
67
-30 300
l 45^
400
500 Temperature in K
Figure 4-52. Disappearance of the amorphous structure of the Mg 72 ^ 2 7 9 alloy during exposure for 13 days at ambient temperature as shown here by successive differential scanning calorimetry (DSC) analysis scans after x hours, as indicated (Hehmann, 1990 a).
186
4 Magnesium-Based Alloys
ticular for amorphous magnesium alloys with acceptable thermal stability (Sec. 4.4.7.1). The constitutional and thermal thresholds associated with the formation of metastable Mg structures were analyzed by one of the authors (Hehmann, 1984, 1988, 1990 b). This topic is reviewed here along with the currently evolving varieties (see also Sec. 4.4.8). 4 A J.I Amorphization by Liquid Quenching
More than a decade ago Sommer et al. (1982 a, b) investigated the extended glass formation ranges formed by ternary additions such as Ag, Zn, Al, Sn, Pb, Sb, and Ca in the Mg-Ni and Mg-Cu binary systems which easily form metallic glass. At that time, research on ternary glass forming Mg systems concentrated on electronic transport and thermodynamic properties associated with the departure from the ideal metallic bonding in magnesium that underlies the resultant mechanical properties, which are explored at a later stage (Hehmann and Jones, 1986, 1987, 1990). Stimulated by the work of Sommer et al. (1982a, b), Hehmann (1985) and co-workers found exceptionally high hardness values for splat-cooled and partially amorphous Mg-Ni and Mg-Ce foils, which then led to the development of prior-art fully and partially amorphous Mg-Ni-Ca foils (Hehmann, 1986; Hehmann et al., 1986, 1988), the latter corresponding to outstanding UTS values of up to 1150 MPa or the highest specific strength values ever reported for a metallic material (i.e. up to 600 MPa cm3/g, see Fig. 4-53). These partially and fully amorphous alloys showed a large ageing response and attractive corrosion properties after selected conditions of heat treatment. Subsequent work by Inoue et al. (1988 a, b, 1989,
1991), Aikawa and Taketani (1991), and Masumoto et al. (1992) confirmed the observations and predictions by Hehmann et al. (1986) for similar ternary systems by adding Sr, Ga, La, Ce, misch metal, and Y to RS Mg-Ni- and Mg-Cu-base alloys. UTS values of the latter ternary and quaternary magnesium-based glasses were above 1000 MPa with the highest specific tensile strength of 436 MPa cm3/g for amorphous Mg 90 Ni 5 La 5 (Fig. 4-53). The amorphous alloys by Inoue etal. (1991) and Masumoto etal. (1992) were more ductile at the expense of specific strength compared to those by Hehmann et al. (1986). The observed coupling factor kH for UTS = (kH x microhardness) ranged from 3.3-3.7, confirming the kH value of 3.5 used for Mg-Ni-Ca (Hehmann et al., 1986). Typical features of such glass forming systems include (i) crystallization doublets at temperatures Tx between 120 and 200 °C upon DSC, indicating their high susceptibility to natural ageing, (ii) pronounced glass transition temperatures, as were evident for more thermally stable glasses with r x >200°C, (iii) relatively large temperature intervals of up to 60 K between the observed glas transition and crystallization temperatures, and (iv) significant softening prior to crystallization at temperatures above 50 °C. The specific hardness data by Hehmann et al. (1986), however, are still 40% higher than those by the Japanese group. Evidently, Ca is an excellent prior-art alloying addition for RS Mg-Ni alloys. The development has since concentrated on amorphous Mg-Ni- and Mg-Cu-base alloys (Inoue et al., 1991; Li et al., 1992b). Many explanations were forwarded to interpret the easy glass formability in magnesium-base alloys. The more advanced interpretations refer to the structural state of the undercooled liquid, including the large
4.4 Rapid Solidification Processing of Wrought Magnesium Alloys
187
600 Mg-Ni-X
Mg-Cu-X I
500
400
|
300
00
'IS 200 O
S
100 0
Figure 4-53. Specific tensile strength values of amorphous magnesium alloys reported between 1977 and 1992. Left panel: Calka et al. (1977); middle panel: Hehmann et al. (1986); right panel: Inoue et al. (1991).
increase in viscosity with increasing undercooling, as is shown by the Vogel-Fulchertype viscosity and which dictates liquid diffusivity instead of the often employed Arrhenius-type diffusivity, and shortrange ordering in the liquid dissimilar to the corresponding equilibrium crystalline long-range order (Sommer, 1993). Using Vogel-Fulcher-type diffusivity in the more recently established growth models by Kurz et al. (1987) and Lipton et al. (1987) shows a dramatic decrease with increasing levels of alloying in the front velocity required to generate massive solidification, as is exemplified for the system Mg-Ca in Fig. 4-54 (Hehmann, 1995 b). On-going development takes advantage of such low solidification velocities for segregationfree solids in the glass-forming concentration range of magnesium alloys. The application of chill-mold-type casting methods
was reported to be sufficient to yield amorphous cross sections of critical thickness d up to 4.0 mm in Mg 80 Y 10 Cu 10 (Inoue et al., 1991) and d— 3.5 mm in amorphous Mg 65 Ni 20 Nd 15 alloys (Li et al., 1992b). It remains to be seen whether these amorphous alloys will find applications in view of their low thermal stability and/or higher densities relative to pure magnesium. Due to alloying, the densities of these alloys are much higher than that of aluminum. 4.4.7.2 Solid State Amorphization Alternative routes towards improved thermal stability of amorphous magnesium alloys include mechanical working of solid magnesium-base mixtures. The potential for amorphization has been demonstrated by cold rolling of alternating Mg-Ni multilayers (Ameur et al., 1991),
4 Magnesium-Based Alloys
188
-175 K
M g - 2 . 0 at.% Ca
1 \
o
J*
600 300
c o
100
CD
c G
10
SUE.
••s CO
o
a>
MS:
1
u
c >. "o
fJCe
PA:
0
-c 3000 2000
oeffi cien
RW:
o
O
-1
|g "o
co Qi
10
X
CO
-2 10 0.0
0.2
0.4
0.6
0.8
1.0
Dimensionless Cross-Section x/d
Dimensionless Cross-Section x/d
Figure 4-54. Internally (Fj) and externally (Fe) controlled growth rates over the dimensionless cross section. External control starts when Ve>V{, the latter taking into account six different heat transfer coefficients h [the numbers on the right given in kW/(m2 K) for rotating-wing (RW) (2-3000), piston-and-anvil (3-600) (PA) splat cooling and melt spinning (10-100) (MS)] assuming Newtonian cooling conditions, as indicated by the horizontal bars. The internally controlled growth rates show dramatic decreases in the solidification front velocity as a function of cross section and level of alloying of calcium in magnesium. An increase in the calcium content by a factor of four reduces the solidification velocities in the instant of initial undercooling by 175 °C from a theoretical velocity of 105 cm/s controlled by thermal diffusion to about 4 cm/s controlled by solute diffusion. The minimum in the internally controlled growth rate after an undercooling of 175°C of the more dilute alloy results from the transition from the thermal to the solute case with a maximum corresponding to the speed of sound. Results are shown for an initial undercooling ranging from 25 to 175 °C with its mean at 100 °C (dashed lines) (Hehmann, 1995 a; after Boettinger et al, 1986).
and by mechanical alloying of pre-alloyed crystalline Mg-54Zn (30 at.% Zn) (Calka and Radlinski, 1989) and Mg-Al-Cabase elemental powders (Hazelton, 1991). More recently, King and Hehmann (1992) reported crystallization peaks of partially amorphous (WE54 + 3-9 mass% A12O3) powders made by high energy ball milling to extend to temperatures up to 500 and 600 °C. The initial strength values of corresponding as-extruded alloys were above 400 MPa at ambient temperatures (elongations of 3-8%) and above 200 MPa at 250 °C (elongations of 50-60%) without having optimized the thermo-mechanical processing. A dispersion of yet unidenti-
fied precipitates of a size smaller than 1 jim was found to suppress recrystallization upon extrusion at 380 °C. It was suggested that exothermic in situ reactions between the alloyed constituents allow for the formation of novel refractory dispersions with more than a doubled volume fraction compared to the unmilled blend of pre-alloyed powders (King and Hehmann, 1992; Hehmann, 1992 b). 4.4.8 Hypersaturated Solid Solutions Passivation by alloying is by definition the improvement in the corrosion behavior of pure base metal by means of alloying.
4.4 Rapid Solidification Processing of Wrought Magnesium Alloys
Ion-implantation of the AZ91 alloy with a dose of 10 17 Fe + ions/cm2 was claimed (Akavipat et al., 1985) to result in passivation. This phenomenon was considered to be related to the homogeneous distribution of the Fe species on the atomic length scale, since Fe inclusions are well known to be one of the most disastrous micro-galvanic prerequisites for catastrophic failure in the history of magnesium alloys. Ionimplanted surfaces, however, are too thin to allow for the formation of Mg surface films with a self-healing capability after being scratched in air. In 1987, Hehmann and co-workers found (Hehmann, 1987, 1988; Hehmann et al., 1989) that extended solid solutions, such as that of aluminum in h.c.p. oc-Mg, exhibit novel anodic polarization features including an anodic plateau at very low current densities, substantial shifts in the pitting potential to more noble values, and annual corrosion rates derived from Tafeltype slopes in aerated aqueous Cl-containing electrolytes of pH less than 5.0 (Fig. 4-55) which were as low as 1 x 10" 3 mm/ year in some cases and ranged up to 0.25 mm/year. By comparison, pure magnesium was observed to show corrosion rates well above 0.25 mm/year under identical experimental conditions. On the basis of these observations, the selective dissolution of either magnesium or deliberately alloyed solutes from corresponding extended solid solutions in oc-Mg was proposed as the principal prior-art mechanism leading to stabilization and improved topological coherency of the surface oxides and hydroxides that form when magnesium and its alloys are exposed to dry air or humid environments. These predictions were based on the observations by Tammann (1919) and Tischer and Gerischer (1958) who reported increasing pitting potential and decreasing anodic current densities of C u -
Log [Current Density
189
(uA/cm2)]
Figure 4-55. Potentiodynamic polarization of singlephase solid solutions of 17.6-23.4% aluminum in ot-magnesium made by melt spinning, and of the chillcast Al 8 Mg 5 phase containing 62.3 % aluminum in a relatively aggressive electrolyte (aerated 0.001 mol NaCl solution of pH4.9 at a temperature of 25 °C) using a scan rate of 6mV/s. The resultant anodic polarization features were not evident for the corresponding two-phase casting materials (Hehmann e t a l , 1989).
Au and Ag-Au solid solutions with increasing levels of the more noble and also passivating alloy constituent, in this case Au. Subsequent work by Pickering and Wagner (1967) gave evidence that the stability of these surface oxides is inversely related to the one-dimensional flux of the volume-controlled diffusion of the less noble constituent through a thin, solid solution surface layer. On this basis, Hehmann (1987, 1988) and Hehmann etal. (1989) showed that this principle also applied to magnesium-base solid solutions, i.e. to magnesium-base alloys with more than 50 at.% magnesium. Accordingly, an aluminum-enriched reaction alloy layer at the surface of the terminal solid solubility extension (TSSE) of aluminum-a-Mg should correspond to a thickness of 15-60 nm, promoting the formation of more stable and more coherent oxide spinels such as MgAl 2 O 4 (Hehmann, 1988; Hehmann
190
4 Magnesium-Based Alloys
etal., 1989; Baliga and Tsakiropoulos, 1992; Baliga et al., 1992). On-going research and development has since added the aspect of passivation by the TSSE of rapidly solidified magnesium alloys as a novel subject to the otherwise enhanced response to ageing afforded by TSSE, being until today solely considered as a precursor phase for dispersion strengthening of advanced magnesium alloys. Both areas are reviewed in the following. 4.4.8.1 Passivation of a-Magnesium An increasing number of research programs are currently confirming the observations and predictions on the microstructural bulk mechanisms proposed by Hehmann et al. (1989) via novel (surface) phenomena associated with TSSE in a-Mg. Makar (1991) and Makar etal. (1992) showed that the repassivation time to establish a current density of 0.1 mA/ cm2 decreased from about 70 s for meltspun Mg-lAl ribbon down to less than 20 s for melt-spun Mg-9A1 ribbon using an aqueous solution of potassium chromate and sodium chloride. In both alloys aluminum was retained in the h.c.p. a-Mg solid solution. Baliga (1991), Baliga and Tsakiropoulos (1992), and Baliga etal. (1992), reported a decrease in the oxide thickness from 100-200 nm on ingot-processed Mg- 16A1 down to 10-50 nm on the corresponding splat-cooled alloy, with aluminum retained in the a-Mg solid solution underlayer. An increasing departure from the equilibrium microstructure of the underlayer (ingot < atomized powder < assplatted) was observed with an increasingly homogeneous distribution of the aluminum cations and the formation of a less permeable MgAl 2 O 4 spinel in the oxide (Baliga, 1991; Baliga and Tsakiropou-
los, 1992; and Baliga et al., 1992), which eventually transformed to hydrotalcite and manasseite (which are derivatives of brucite) with an Mg 2 + /Al 3 + ratio of 3:1 (Baliga and Tsakiropoulos, 1993). A number of trivalent solutes were proposed as candidates to uniformly substitute for magnesium and/or aluminum in the surface oxide, if a sufficient departure from microstructural equilibrium was achieved and if the corresponding solute has an ionic radius similar to Mg 2 + and Al 3 + (e.g., P, Ga, Co, Cr, Fe, Mn, Ni, V, and Ti (Baliga and Tsakiropoulos, 1992). An enhanced response to anodic polarization was also reported for the solid solution of 2.7% Zr in a-Mg made by laser cladding, and resulted in an improved resistance to corrosion compared to alloy AZ91B (Subramanian et al., 1991). On the basis of these observations it appears straightforward to observe such improvements for passivating components other than aluminum in magnesium. More recent research and development programs concentrate on methods in which the alloying of magnesium is not limited by liquid immiscibility with relevant constituents. This topic is the subject of another article in conjunction with forecasting applications of such structures that have. yet to be established (Thorbeck and Hehmann, 1996). The best yet confirmation of the hypotheses that selective dissolution in TSSE of a-Mg leads to solute cation enrichment and stabilization of the surface oxide (see above) was put forward by Hirota et al. (1993), who reported on anodic polarization in 1 mol HC1 aqueous solution for the extended solid solution of 8-47 Ti, 20-77 Zr, 14-72 Nb, and 18-77 Ta (all in at. %) in h.c.p. a-Mg made by DC magnetron sputtering (Fig. 4-56). Substantial transition metal (TM) cation and corresponding O 2 ~-anion enrichment and
4.4 Rapid Solidification Processing of Wrought Magnesium Alloys (a)
3. .•& c
4.4.8.2 Novel Constitutional Thresholds for Enhanced Response to Ageing
\47Ti
Mg
o
1 Ti 0
-1
Ti
•
c 3
Mg - xTi
-2-
1n HCI 3CTC
3
191
-3
-2
-1
0 SCE Potential in V
1
-0.5 0 SCE Potential in V
0.5
(b)
-1.5
Figure 4-56. As for Fig. 4-55; here for single-phase solid solutions of (a) titanium and (b) zirconium in a-magnesium made by magnetron sputtering and using a 1 mol NaCl electrolyte at pH 9 and 30 °C. The effect of this electrolyte on the pure elements is also shown (Hirota et al., 1993).
Mg 2+ -cation and OH~-anion depletion was observed to underlie the passivation in this electrolyte at applied potentials above -0.68 V (Mg-47Ti), - 0 . 3 V (Mg-57Nb), and -0.25 V (Mg-38Ta) with respect to the saturated calomel electrode potential (Fig. 4-57). The thickness of air-formed oxides was 2 nm, while that formed in the electrolyte was not more than 4.5 nm. Annual corrosion rates for solute concentrations of more than 40 at.% decreased from 0.03 mm/year for the Mg-Ti system down to as low as 1.5 x 10" 3 mm/year for the Mg-Ta system. Hirota et al. (1993) considered the crystalline state of the extended solid solutions to prevent sufficient passivation.
The number of elements required to form an extended solid solution in oc-Mg by nonequilibrium solidification processing has increased from 5 elements in 1986 to 36 elements in 1993 (Thorbeck and Hehmann, 1996). The maximum equilibrium solid solubility was extended by factors ranging from 1.5 for Ag to about 1000 for Ba in magnesium by quenching from the melt (Hehmann, 1990b; Hehmann et al., 1990). This new development allows for a number of novel applications (Thorbeck and Hehmann, 1996), though no approach has yet been reported to systematically explore TSSE. The transformation sequence of the extended solid solution of Y in oc-Mg involves 11 partially exothermal or endothermal microstructural evolutions (Sommer et al., 1990), while that of aluminum in oc-Mg involves 6 such reactions, as identified by TEM and DSC work (Hehmann, 1988, 1990 a). These developments have led to significant progress in the state-of-the-art knowledge of the precipitation kinetics of magnesium alloys, though not yet to everyone's satisfaction (see Polmear, 1992). The extension to up to 22% Al in oc-Mg made it possible to confirm (Hehmann, 1990 a) an as yet unidentified y'-phase that forms upon natural ageing prior to endothermal transformation and exothermal decomposition of the corresponding solid solutions at higher temperatures (Fig. 4-58). Until recently, the interpretation was that the transformation of the solid solution of aluminum in oc-Mg leds directly to the formation of the equilibrium two-phase mixture (a + y) where y = Mg 17 Al 12 (Clark, 1962). Obviously, aluminum levels below the equilibrium terminal limit do not result in a sufficient volume fraction for the identification
192
4 Magnesium-Based Alloys
0.8
5 0.6
c 0.8 o "4 \ 0.6
Mg - 47Ti 1n HCI30°C
o
o 0.4 03
O
0.2
0.2-
>
Nb 5+
;
Mg-57Nb
\> '\
1
4
1n HCI30°C \ \
Mg 2 + t
As P.
Ec-0.6
-0.4 -0,2 Potential in V
0.5
0
AsS.
I Mg - 47Ti 1n HCI30°C
As P.
Er-0.6
-0.4
1.5
2 Mg -57Nb
LL O
I 1
-
1n HCI30°C
W A
-0.2
Potential in V
1 Potential in V
0.5
AsS.
1 Potential in V
1.5
Figure 4-57. Cationic and anionic fraction in the surface oxide of hypersaturated solid solutions of 47 at.% (63.6%) titanium or 57 at.% (83.5%) niobium in a-magnesium upon polarization, as in Fig. 4-56. "As P.": as prepared; "as S": surface film formed upon exposure air after sputtering. Ec is the corrosion potential (Hirota et al, 1993).
of the y'-superlattice by electron microscopy because of the similar atomic structure factors for solute aluminum and solvent magnesium (Clark, 1986). Rohklin et al. (1990) suggested that the absence of Sn-containing Guinier-Preston zones and of equilibrium Mg2Sn precipitates resulted in the low age hardening response of + 1 0 % which was observed for solution treated Mg-9.9 %Sn. Vostry et al. (1991) observed an increment in the microhardness of about 40% upon ageing binary Mg-10%Ca melt-spun ribbon, while a response of +100% was found for melt-spun Mg-15%Y ribbon (Vostry et al., 1992). Most of these observations relate to a quasi one-dimensional extension of the solid solubility in oc-Mg as a function of
solute concentration, while the front velocity and corresponding supersaturation slow down dramatically upon traversing the particulate and foil cross section (see Fig. 4-54) due to the evolution of latent heat that is dissipated back into the recalesced, nonsolidified liquid part of the fragmented alloy volume. The heat flow is then controlled by the dependence of the specific alloy heat on undercooling. More recent effort was therefore devoted to investigating the limitations imposed by the kinetics for liquid solidification more closely (Miiller et al., 1992; Hehmann and Tskiropoulos, 1992; Rohklin et al., 1991). Miiller et al. (1992) reported a maximum width of about 20 \xm for extended solid solutions in a-Mg by combining a heat flux model for external heat flow with the mod-
4.4 Rapid Solidification Processing of Wrought Magnesium Alloys
193
maxima are a result of competition between the two principal growth modes for segregation-free solidification, i.e. absolute stability and solute trapping. Hypoeutectic Mg-Sr alloys, for example, impose the need to employ high speed lasers with rotating sample cylinders in order to obtain corrosion properties that are reliable for weathered applications.
(a)
4.4.9 Development Principles for Rapid Solidification Processed (RSP) Wrought Magnesium Alloys (b)
4.4.9.1 Damage Tolerance by "Clean" Grain Refinement
300
400
500
600
700
Temperature in K
Figure 4-58. (a) TEM diffraction pattern and (b) DSC study using various heating rates, as indicated in K/s for melt-spun Mg-23.4A1 ribbon after 12 months exposure to ambient temperatures. The endothermal transformation of the ordered room temperature phase is as strong as the transformation of the remaining supersaturated solid solution of aluminum in a-magnesiumn after these conditions of natural ageing (Hehmann, 1990 a).
els for dendritic and planar crystal growth, as well as for the specific heat of relevant magnesium alloy melts. Hehmann and Tsakiropoulos (1992) showed that front velocities required for segregation-free solidification of hypo-eutectic Mg-Sr alloys can extend to a maximum of 4 m/s under conditions of externally controlled heat flow, so corresponding to laser withdrawal velocities of up to 6 m/s and more. Such
One of the prime objectives of micro-alloying by conventional processing is microstructural refinement in order to increase the strength, thermal stability, and deformability, as well as to reduce microshrinkage of the corresponding magnesium alloys. Many elements considered for micro-alloying (Pekgiileryuz et al.? 1992, 1993) are low melting point trace elements which might recontaminate the recently established HP ingot magnesium alloys (see Sec. 4.3.1.1) by forming micro-galvanically active second phases. Moreover, microalloyed constituents have a tendency to be redistributed on grain boundaries when processed by conventional casting methods as a result of partition coefficients below unity. It should be noted that many corrosion failures in the history of magnesium alloys occurred in real life service as a result of micro-galvanically active inclusions. Microstructural refinement by rapid solidification ( < 1 }im; Sec. 4.4.2) does not only result in more effective grain refinement than micro-alloying by conventional means (10-30 jam; see Sees. 4.3.1, 4.4.9.2, 4.4.9.3, and Fig. 4-59). It also has the advantage of higher chemical homogeneity,
194
4 Magnesium-Based Alloys
5 I
50 10
i f
2 I
1 I
0.5 1
0.3 1
0.2 1
500 |Mg-6Zn-0.45Zr| 400
y
300
-
200
/A
100
/
0.0
|
RSAZ91
y r^ | Mg - 9 to 13 Al base jg
/Ay^A QE22 \ Chill-Cast AZ91 || | pureMg ||
m
V/ /
|
|(a+p)Mg-9Libase|
| bcc-p Mg - 40Li - 2H j | 1
1
1
0.5
1.0
1.5
1
2.0
2.5
Figure 4-59. Tensile yield strength ay as a function of d~1/2 (d is the grain size) for extruded magnesium alloys made by ingot processing and by rapid solidification processing. From Hehmann (1995 b), here without individual samples indicated. Note that the minimum grain size observed is represented by the end on the right-hand side of the Hall-Petch relationship.
so avoiding undesirable and unpredictable micro-galvanic effects which may be introduced by conventional micro-alloying. Cotton and Jones (1989, 1991), for example, reported one to two orders of magnitude lower corrosion rates for RS M g 15%A1 splats compared to I/M AZ91 as a function of the Fe impurity level. This improvement was considered to stem from matrix ennoblement with aluminum (Sec. 4.4.8.1) and a reduced rate of proton discharge around the refined Fe inclusions. While conventional micro-alloying might increase the susceptibility of HP Mg alloys to corrosion, RSP increases the damage tolerance of nonHP Mg alloys and would improve the corrosion resistance of HP Mg alloys even further.
Under compressive loading and under tensile loading at hydrostatic pressures of 230-700 MPa RS AZ91 was found (Lahaie et al., 1992) to show macroscopic and localized shear bands without premature grain boundary fracture and without any evidence of twinning. The grains were 1.2 + 0.5 (im in size and decorated by Mg 17 Al 12 particles of 0.1-0.3 jim in size. I/M AZ91 of grain size 8-15 |im, however, showed ample evidence of twinning under such conditions. The observed shape change and the resulting microscopic strain of the 1 jim grains were smaller than the imposed macroscopic strain. At room temperature using a strain rate 8 = 1 0 ~ 4 s - 1 (Lahaie et al., 1992), the true (engineering) strain-to-fracture of RS AZ91 increased from 0.1 (10.5%) for a hydrostatic pressure of 0 MPa up to 1.6 (400%) for a hydrostatic pressure of 700 MPa. Grain refinement down to about 1 jim was therefore considered to activate new deformation mechanisms which allow grain boundary sliding and new flow processes at ambient temperatures, so substantially improving the ductility of wrought magnesium alloys. These results not only confirm the ductility transition in the grain size regime of 5-10 jam (Fig. 4-7), as was shown by Emley (1966) for pure magnesium (Sec. 4.2.2.3). They also suggest that such grain refinement is difficult to achieve without using RSP. 4.4.9.2 Effect of Grain Refinement on Hall-Petch Strengthening The compressive stress to fracture of RS AZ91 was between 650 and 750 MPa at T= 77-273 K and a strain rate s of 10~ 4 s" 1 . At room temperature using a strain rate s = 10 " 4 s ~* (Lahaie et al., 1992), the tensile stress to fracture of RS AZ91 increased from 400 MPa for a hy-
4.4 Rapid Solidification Processing of Wrought Magnesium Alloys
drostatic pressure of 0 MPa to up to more than 900 MPa for a hydrostatic pressure of 700 MPa. The latter strength value corresponds to a value of E/70. Localized shear at stresses of the order of the theoretical shear strength was considered to become the limiting strengthening mechanism if intragranular grain boundary failure was suppressed via grain refinement. The validity of such postulations can be examined by the response of the HallPetch proportionality constant ky to grain refinement. Nussbaum et al. (1989 b) established a Hall-Petch-type relationship for the yield strength (TYS) of the unmodified RS AZ91 composition as a function of the grain size d TYS = ao + kyd~112
(4-3)
with <70 = 130 MPa and the locking coefficient ky = 210 MPajum 1/2 . Jones (1991) suggested that the TYS data for Si- and RE-containing RS Mg-Al-Zn-based alloys developed by Allied Signal are in agreement with these values by Nussbaum et al. (1989 b). In fact, the values by Allied indicated somewhat higher <J0 and ky values (see Hehmann, 1995 b, and Fig. 4-59). Evidently, Si and RE act not only as the often quoted microstructural stabilizers via the formation of temperature stable dispersoids that pin the motion of dislocations and grain boundaries upon consolidation of pulverized ribbons, flakes, and droplets, but also as nucleants and catalysts for more effective grain refinement itself during both solidification and/or consolidation. The first and basic effect of such additions is thus to increase the boundary strengthening effect afforded by grain refinement upon the entire processing route compared to those alloys without such additions. Already as long as twenty years ago, Isserow and Rizzitano (1974) reported on
195
grain refinement by RS as one of the dominant factors for improving the strength of the ZK60 alloy, using "micro-quenching". The corresponding ky value was in the region of that reported by Nussbaum et al. (1989 b), while the intercept a0, however, appeared to be significantly higher than for Mg-Al-Zn alloys (Fig. 4-59). A ky value of 185 MPa jim1/2 was evident from the literature on the magnesium alloy QE22 (Vervoort and Duszczyk, 1991). Meschter (1987) reported a strength increment of 68 MPa for solution treated RS Mg-10.9Al over ingot processed Mg-9A1, which was consistent with a grain size reduction from about 100 to 11 |im, i.e. by a factor of 10, assuming a Hall-Petch factor ky of 337 MPa |im 1/2 . While the ky values for the above precipitation-hardened Mg-Al-Zn-based and ZK60 alloys are significantly lower compared to pure magnesium (see Fig. 4-59; Cottrell, 1975), the value for the monophase Mg-Al alloy is above that for pure magnesium (280 MPa jim1/2) and conventional magnesium alloys (280 MPa ^im1/2; Hall, 1970). The data Meschter (1987) reported for the multiphase RS Mg-13Al-1.2Si alloy are in excellent agreement with the values for binary Mg-lOAl alloys, as extrapolation to smaller grain sizes shows (Fig. 4-59). Finally, the reported data for Mg-Li based alloys result in 100 MPa jum1/2 for the (oc + P) Mg-9Li-base alloys (Meschter and O'Neal, 1984; Meschter 1986) and in 58 MPa [im1/2 for b.c.c. P-Mg-6Li-0.12H (Haferkamp et al., 1990; Fig. 4-59). These are the lowest ky data ever observed for a magnesium alloy, resulting in softer alloys than pure magnesium for grain sizes below 40-50 jam. Microhardness studies confirmed the above positive but small ky values for the (oc + P) and the b.c.c. P-Mg-Li alloys. The data reported for binary h.c.p. oc-Mg and
196
4 Magnesium-Based Alloys
b.c.c. P-Mg-Li alloys (Schemme and Hornbogen, 1992; Schemme, 1993 a), and for the ternary engineering alloy LAI41 (Das et al., 1988) resulted in a ky value of about 50 VHN |im 1/2 (Fig. 4-51). The results show that grain size is the dominating strengthening factor in any of the reported Mg-Li alloys, and that the level of alloying for a given crystal structure does not significantly affect the proportionality constant ky. The intercept a0 of these solid solution alloys depends on both the crystal structure and the alloy system (Fig. 4-51; Hehmann, 1995 b). Li and Jones (1992) observed that the microhardness of EA55RS also follows a Hall-Petch-type relationship with a ky value of 110 + 30 VHN jim1/2, as was confirmed by a decrease in the microhardness and by grain growth on heat-treatments above 300 °C. Microhardness, however, was also reported (Project MAIE, 1990) to scale with volume fraction and decreasing size of the second-phase particle dispersion in RS magnesium alloys. Microhardness studies may therefore be a useful method for discriminating the boundary strengthening effect from other mechanisms like precipitation hardening in multiphase engineering alloys, and to study the rate-controlling mechanism for boundary strengthening over a more detailed regime of grain size (Hehmann, 1995 b). 4.4.9.3 Multiphase Dispersions A common feature of this second generation of engineering RSP magnesium alloys is their high volume fraction of precipitates and/or dispersions concentrating on cell and grain boundaries and/or traversing the refined grains homogeneously (Fig. 4-60). Juarez-Islas (1991) reported on Mg 1 7 Al 1 2 , Mg 7 Zn 3 , and Mg2Si phases concentrated on the cell boundaries. These
(a)
20 pm
(b)
Figure 4-60. (a) Optical and (b) TEM micrographs of RS AZ91 + 2 % Ca after extrusion at 250 °C using an extrusion rate of 20:1. The TEM micrograph shows Al2Ca-type precipitates of size 10-70 nm (courtesy of Pechiney).
were already present in the as-solidified, splat-cooled condition of Mg-12A12.5Zn-0.5-3.0Si alloys. The overall appearance of the as-extruded microstructure of the new alloys of Pechiney and Norsk was considered to be independent of compositional details (Nussbaum et al., 1991). These alloys have equiaxed oc-Mg grains of size 0.3-0.5 jim which are slightly supersaturated with aluminum (Nussbaum et al., 1990a, 1990b, 1991), and a homogeneous dispersion of second phases of size 0.01-1.0 jim. The smaller type (Mg2Si, Al^RE^, AlaAEb) are transgranular, while the larger Mg 17 Al 12 type are concentrated on grain boundaries. The Mg17Al12-type phase was reported to contain Ca and Zn in solid solution and
4.4 Rapid Solidification Processing of Wrought Magnesium Alloys
to transform to the binary equilibrium Mg 17 Al 12 phase upon heat treatment, without dissolution. The observed differences in the effect of Ca and Sr upon the corrosion behavior and ductility of RS AZ91 could result from the different chemical compositions and the volume fraction of the resultant second phase dispersion (Sec. 4.4.3.1 and Table 4-24). The second phase dispersion of the RS AZ91 + 2 % Sr alloy was reported to consist of Mg 17 Al 12 and eventually Mg2Sr and Al4Sr, while MgflCab-type second phases were not observed in the RS AZ91 + 2 % Ca alloy. A metastable Ca- and Al-containing intermetallic phase was found instead which transformed to Al2Ca of size 0.05 |im upon heat treatment (Gjestland et al., 1991; Nussbaum et al., 1991). About 1011 particles/mm3 were estimated to be homogeneously distributed throughout grains of size 0.6 |im in the RS AZ91 +2.3% Ca alloy (Gjestland et al., 1991). They were considered to pin dislocations and grain boundaries more effectively than the remaining phases (Sec. 4.4.3.2). RE- and AEcontaining second phases tend to be less noble than the magnesium matrix and could upon precipitation reduce the effective anodic area exposed to the environment (Hehmann, 1988). Similar observations were made for heat-treated M g - C a Cu splats showing dramatically improved corrosion resistance upon precipitation of Mg2Ca from the corresponding supersaturated solid solution (Hehmann et al, 1986). Such precipitates were considered to neutralize even the very harmful micro-galvanic effects triggered by a coexisting Mg2Cu phase. The Al2Ca phase might play a similar role to Mg 2 Ca. In contrast to previous reports (Das and Chang, 1989; Chang et al., 1987), Li and Jones (1992) did not find any Mg 17 Al 12 or Mg 3 Nd phases in the microstructure of the
197
Allied EA55RS alloy. It should be noted that the Allied alloys have the highest Zn content of the engineering RS magnesium alloys reviewed here (around 5%). Two microstructural zones, i.e. fine and elongated oc-Mg grains of diameter 0.5 Jim along the traverse section and coarse grains of up to above 5.0 jim, were found to coexist in the extruded EA55RS alloy. As-extruded samples and samples isochronally heat treated for 1 h at 185 °C showed a large volume fraction of fine (i.e. 20 nm) and globular Al 2 Nd phases and relatively coarse (i.e. 0.1-0.2 |im) icosahedral ternary (Mg-Al-Zn) phases. Both phases delineate primarily the boundaries of grains and subcells in both types of microstructure. The size of the Al 2 Nd phase did not change on heat treatments up to 400 °C. However, the icosahedral phase dissolved at temperatures above 300 °C. The Al 2 Nd phase was considered to form during rapid solidification, while the icosahedral phase was considered to correspond to the temperature instable icosahedral Mg 32 (Al,Zn) 49 phase previously obtained by RSP and mechanical alloying of that composition itself (Ivanov et al., 1987; Graebner and Chen, 1987). The results were consistent with those of Regazzoni et al. (1991), who found that the Mg 17 Al 12 phase is replaced by a crystalline Mg 32 (Al,Zn) 49 phase in RS Mg-Al-based alloys with Zn levels of more than about 2%. A similar phase forming ability in more dilute magnesium alloys, however, remains to be identified. At temperatures above 300 °C, an Al 11 Nd 3 related phase containing Zn was observed at the boundaries of the smaller grains, while coarse MgZn precipitates were also found at high angle grain boundaries (Li et al., 1992 a). The Al x l Nd 3 related phase (r m = 1235°C) was somewhat larger (i.e. 70 nm) than the coexisting
198
4 Magnesium-Based Alloys
Al 2 Nd phase (Tm = 1460 °C). The resultant second phase dispersion delineates the cell boundaries which are not entirely uniformly distributed. However, EA55RS has more stable compounds than are possible by conventional processing such as Mg2Si (r m = 1085°C). Corresponding "RSP"phases were at least one order of magnitude smaller than was reported for I/M processed magnesium alloys (Sec. 4.3.2.2). Both factors contributed to the resistance to degradation of the microstructural refinement obtained by RSP, as is necessary for Mg-Al-Zn-alloys upon consolidation and at higher service temperatures.
4.5 Magnesium Metal Matrix Composites Numerous applications of lightweight materials require high stiffness and low thermal expansion. Unfortunately the stiffness and thermal expansion of magnesium are inferior to competing materials, and substantial improvements are not available by conventional means. Magnesium metal matrix composites (MMCs) offer one solution to these drawbacks. Materials with high elastic moduli and low thermal expansion such as oxides, ceramics, graphite, intermetallics, etc. (see, e.g. Vol. 13 of this Series) are used to overcome these limitations in a way that is ideally predictable by the rule of mixtures (ROM), i.e. the additive behavior in relation to the elasticity and the thermal expansion of the magnesium matrix and the employed reinforcement. The resultant improvements depend, however, not only on the properties and the volume fraction of the reinforcement, but also on the shape (long fibers, short fibers, particles, plates) and the resultant interface between the reinforcement and the magnesium matrix. Other proper-
ties of magnesium to be improved by such reinforcements include wear resistance, tensile properties particularly at elevated temperatures and creep resistance. These advantages are usually obtainable at the expense of fracture strain, machinability, and corrosion resistance. The variety of processes for producing magnesium MMCs include all of the techniques used to produce aluminum MMCs (see, e.g. Vol. 13). Processes used or under investigation are the stirring of particles in the melt followed by casting the resultant mixture, liquid metal infiltration (squeeze casting), self-generated vacuum infiltration, powder metallurgy and extrusion, the supplementary use of PVD and CVD techniques or diffusion bonding, and more recently also co-spray forming techniques. As is the case in aluminum MMC technology, the production of magnesiumbased MMCs with unidirectional continuous fibers requires the use of PVD, CVD, or diffusion bonding techniques. More readily available production techniques like liquid metal infiltration cannot be used to manufacture equidistant fiber arrangements in the alloy matrix. This is a particular problem for magnesium MMCs with a low fraction on fibers. A high reactivity between the magnesium and the reinforcement results in better interfaces than by using an aluminum alloy matrix. In several magnesium matrix-reinforcement combinations, however, this reactivity leads to limitations in elevated temperature applications due to severe interface reactions. Many materials have been used or tested to reinforce magnesium. The most important are graphite fibers, SiC particles, and A12O3 fibers and particles. Since magnesium does not form carbides easily, carbon fibers are an ideal reinforcing material for magnesium. The low density of graphite
4.5 Magnesium Metal Matrix Composites
fibers gives a composite material with a density of less than 2000 kg/m3. A lot of attempts have been made to coat graphite fibers in order to improve wetting. Composites made with these fibers, however, did not show the expected good properties (for details, e.g. see the review by Pekguleryuz, 1991). A variety of experiments have been performed using silica-coated graphite fibers in order to improve wetting. Chin and Nunes (1988) reported superior longitudinal and transversal mechanical properties for Gr/AZ91C compared to Gr/Mg and Gr/ZE41A by using an amorphous SiO2 coating on these fibers. Fiber-matrix degradation was observed due to a reaction between the coating and the rare earth elements. This investigation showed the importance of selecting the proper fiber-matrix system. Diwanji and Hall (1987) discovered that uncoated and untreated graphite fibers in Mg-Al alloys resulted in about two times better mechanical properties than with surface-treated fibers. The bonding between the matrix and the fibers was good. The aluminum in the matrix reacted with the fiber surface to form fine A14C3 precipitates, and this reaction zone was not affected by heat treatments. Attractive mechanical properties were achieved by squeeze casting without coating the fibers (Hall, 1991). Surface treatment improved the interfacial shear strength, but reduced the tensile strength. Alloying elements improved the tensile properties provided that no continuous crack paths occurred through the resultant brittle phases (Hall, 1991). Bode and Hornbogen (1992) reported good interfacial bonding between graphite fibers and an Mg-1 %A1 alloy. In Mg-Al alloys with high aluminum levels, e.g. AZ91, a rapid formation of A14C3 precipi-
199
tation was observed, thus excluding such alloys from further consideration as a matrix material for graphite reinforcement. Viala et al. (1991) investigated the effect of magnesium on composition, microstructure, and mechanical properties of carbon fibers. The fibers were isothermally heat treated (450-700 °C) under a saturated vapor pressure of magnesium. They concluded that highly graphitized fibers, e.g. pitch-based P55 or P100, remain unaffected by long-time annealing in the presence of magnesium vapor, whereas impure and disordered fibers, e.g. PAN-based T300, undergo some chemical and microstructural modifications which lead to the degradation of their mechanical properties. Tensile properties as high as 1400 MPa and longitudinal Young's modulus of 200 GPa have been obtained with the M40J/AZ61 unidirectional composite (Rabinovitch et al., 1992). Kleine and Dudek (1992) investigated the interface of SiC-coated carbon fibers and did not find any degradation of the fiber nor of the SiC layer. The only reaction product they found was some single MgO particles. Carbon fibers have negative thermal expansion in the longitudinal direction. They are used for composites with zero thermal expansion along the direction of fiber alignment (Fig. 4-61). The difference in the thermal expansion of carbon fibers and magnesium alloys results in large internal stresses being generated during processing and, for example, by thermal cycles upon service. Most of the internal stress is reduced by interface slip. This relaxation is greatly influenced by the structure of the fiber-matrix interface. Armstrong etal. (1990) measured an internal stress of 25 MPa in a Gr/Mg-Zr MMC which is smaller than in respective Gr/Al composites. Dries and Tompkins (1988) reported indications of a thermal strain hys-
200
4 Magnesium-Based Alloys
30
/
-
^
— C2
EZ41
LU
EZ41 + 19vol.-% Saffil® _-— --.
20'.
QE22 + 21vol.-% Saffil®
10
Mg + 67vol. -% unid. C-fibers 7~7TTT-T -Till-*-
0
100
200
300
.
1^—f-r-T"
400
500
Temperature in °C Figure 4-61. Temperature dependence of the coefficient of thermal expansion for EZ41, EZ41 + 19 vol.% Saffil, QE22 + 21 vol.% Saffil, and pure Mg + 67 vol.% unidirectionally arranged C fibers (courtesy of Metallgesellschaft).
teresis for Gr/Mg composites, which was not reduced by heat treatment-based processing or by strain hardening during thermal cycling. Magnesium MMCs with graphite fibers suffer from severe galvanic corrosion even in an ordinary laboratory atmosphere because of the electrically conductive graphite fibers (Hall, 1987). Therefore applications were initially considered to be limited to satellite parts (Armstrong, 1979). Goddard et al. (1986), however, reported on AZ91C/P55 (coated) MMC for lightweight automotive rotary engine housings with good mechanical properties (Table 429), but the corrosion behavior was not included. Most of the carbides are well suited for reinforcing magnesium alloys. Particulate SiC or B4C are relatively cheap and light with densities of 3.1 and 2.5 g/cm3, respec-
tively. These carbides bond well to the magnesium matrix. These magnesium MMCs can be produced by the melt stirring method, which requires particles with good stability against chemical attack by the magnesium melt. Laurent et al. (1992) reduced the reaction between SiC particles and liquid AZ91D by adding SiC at temperatures between the solidus and the liquidus, followed by stirring this mixture above the liquidus temperature. Lim and Choh (1991) reported that the incorporation time of SiC into the molten, pure magnesium was remarkably shorter, and the condition of the homogeneous particulate distribution was better than that of pure aluminum. The incorporation time was not affected by adding Ce, Mn, Zr, Bi, Pb, Sn, and Zn to magnesium, and it was prolonged by adding Al, Cu, Ca, and Si. Ca additions were found to lead to particulate agglomerations. B4C and SiC particulate-reinforced magnesium composites have been made on a pilot plant production scale (Mikucki et al., 1990; Wilks and King, 1992) resulting in a 50 % improvement in the modulus and elevated temperature strength (Table 4-29; see also Sec. 4.4.4.1). The strength of such composites, however, decreases substantially with increasing temperature (Mikucki et al., 1990). A 100% improvement in creep strength at 150°C was reported by Wilks and King (1992) for a ZC71 alloy containing 12 vol.% SiC. The choice of the matrix alloy is important if a sufficient reinforcing effect is to be obtained. The reinforcing particles may influence the age hardening behavior of the matrix. Badini et al. (1992), for example, observed that Mg 17 Al 12 precipitates nucleated heterogeneously close to the interface between the B4C particles and an AZ80A matrix. The ageing response of this MMC is lower than in the corresponding unreinforced alloy.
4.5 Magnesium Metal Matrix Composites
SiC whiskers reinforce magnesium more effectively than SiC particles. However, it is difficult to obtain a homogeneous distribution of whiskers in magnesium and aluminum. Liquid metal infiltration of preforms made from SiC whiskers was often observed to result in agglomerations of the whiskers. The powder metallurgical route leads to severe damage of the fibers during extrusion (Kainer and Schroder, 1990). Both of these effects cause a decrease in the mechanical properties. The risk of inhalation of whiskers by the operator imposes another obstacle to this particular MMC route. SiC whiskers, however, offer the only feasible solution to reinforcing Mg-Li alloys. Other fibers and particles degrade during fabrication and heat treatment (Mason et a l , 1989). Kieschke et al. (1991) developed sputter-deposited coatings for SiC monofilaments. They found that very thin layers of Y 2 O 3 offered some protection, but much development work has yet to be done in order to find an economically viable solution to reinforcing Mg-Li alloys. Much work has already been done on reinforcing magnesium with A12O3 fibers as well as with A12O3 particles. Previous work on A1/A12O3 composites has led to numerous commercially available A12O3 fiber preforms. In contrast to aluminum matrices, magnesium and magnesium alloys show excellent wetting with A1 2 O 3 . From the thermodynamical point of view A12O3 is not stable against chemical reaction with magnesium, resulting in the formation of an MgAl 2 O 4 spinel or MgO. Magnesium reacts even more strongly with silica, which is used as an inorganic binder in commercially available A12O3 fiber preforms, and is present in alumo-silicate-, mullite- and 5-type A12O3 fibers; in the latter to stabilize the 5-phase structure (e.g. Saffil fibers, supplied by ICI). Intermetal-
201
Figure4-62. Microstructure of a ZE41 + 19vol.% Saffil composite, T4, etched: (a) intermetallic phase precipitates on the fiber surface (arrows), (b) grain size in the nonreinforced part of the casting, and (c) grain size in the fiber-reinforced part of the casting (courtesy of Metallgesellschaft).
lie phase precipitates tend to nucleate on fibers and/or the inorganic binder (Fig. 4-62; see also Kainer, 1991 b). The production of Al2O3-magnesium composites requires short exposure of the fibers to the
202
4 Magnesium-Based Alloys
magnesium melt (e.g. by the squeeze-casting process), and the service temperature is somewhat limited. The degree of the fibermatrix reaction depends on the service temperature, the alloying additions to magnesium, and the type of A12O3 reinforcement. Aluminosilicate fibers (50% A1 2 O 3 , 50% SiO2) react completely with magnesium. Highly reactive alloying elements, e.g. REs or Zr, tend to segregate at the reinforcement-matrix interface and so proper selection of the matrix-reinforcement combination is required (Kainer 1991 b). Figure 4-62 shows that this segregation reduces the corresponding grain refining effect of these elements. Kainer (1991b) found that both QE22 (Mg-2.5Ag2.5Nd) and Mg-2.5Nd alloys reinforced with 20 vol.% Saffil fibers also show good UTS properties at RT and at 200 °C (Table 4-29). Respective composites with Mg-2.5Ag and pure Mg matrices lost strength at 200 °C. The segregation of Nd into the fiber-matrix interface was considered to prevent the fiber from severe reaction with the matrix. Ag as an alloying addition in magnesium has the disadvantage of reducing the corrosion resistance. Kainer (1992) showed that the UTS of T6 heat-treated AZ91/Saffil-A12O3 composites is on the same level as for the unreinforced alloy, whereas in the as-cast condition the fiber-reinforced AZ91 has the better values. Sayashi et al. (1988) produced completely A12O3 short fiber reinforced pistons with an AZ91 matrix. This composite part lost strength at elevated temperatures. Hino et al. (1990) achieved better results with QE22 and employed finally a binary M g Nd alloy matrix. Hino et al. (1990) found that the SiO2 binder used to fix the fibers in the preforms, reacted with the magnesium to form Mg2Si particles, causing crack formation and failure.
Schroder and Kainer (1991) made hybrid magnesium composites by infiltration of a preform containing A12O3 Saffil fibers and SiC particles in QE22. The best results were achieved using 10 vol.% fibers and 15 vol.% SiC particles (Table 4-29). This combination gives excellent stiffness, while the high temperature tensile properties are somewhat reduced compared to fiber reinforcement alone. In summary, all magnesium-based MMCs require tailoring of the corresponding interfaces between reinforcement and matrix in order to meet the requirements set by the customer (see also Vol. 13 of this Series). Selection of the magnesium MMC concerned depends primarily on the specific application and temperature in service. This is a large field currently under development for better magnesium-based composites.
4.6 Recycling of Magnesium Alloys The production costs of primary magnesium represent one driving force for building a "secondary" magnesium metal market, as is well established for aluminum. In the 1970s magnesium prices increased rapidly. In 1971 the use of magnesium for the "beetle" peaked at 42000 t/year. Volkswagen reduced the costs of the final magnesium components by using secondary magnesium from internal and external sources (Hollrigl-Rosta et al., 1980). The use of secondary magnesium, however, did not stimulate the use of magnesium at that time. Evidently, the small market for structural magnesium has not yet caused a secondary magnesium market to build up and current magnesium scrap is easily absorbed by the need to desulfurize steel, for example.
4.7 Conclusions and Outlook
Environmental legislation and increasing waste removal and deposition costs have reportedly imposed pressure on industry to establish closed material circuits. As with other metals, the inherent recyclability of magnesium is expected to attract increasing attention. Highly engineered plastics, including fiber-reinforced plastics or electrically conducting plastics, are only recyclable in a "downstream" manner such as by burning. When both recycling and density considerations become of prime importance, any progress in the quality of secondary magnesium-based structural products will provide an undue advantage over secondary aluminumbased structural applications. The image and quality of secondary magnesium has to be improved if it is to become a viable secondary source of prime relevance. This includes the early separation of magnesium alloys from different materials after use, followed by the separation of one magnesium alloy family from another (see, e.g. Sattler and Yoshida, 1992; Bolstad and Roenhug, 1992), as well as the need to establish a recycling trademark for magnesium (Cutler, 1992). Magnesium suppliers offer sufficient recycling capacity and satisfactory technology to customers who cannot afford inhouse recycling and/or recycling of high purity magnesium alloys without undue contamination (0ymo et al., 1992; Housh and Petrovich, 1992). Recent developments include the use of protective gas atmospheres including CO 2 /dry air/SF 6 mixtures to circumvent the use of fluxes. 0ymo et al. (1992), for example, reported a continuously operating, two-furnace system in which the remolten scrap is transferred to a casting unit allowing for melt treatments to result in high purity quality, as for primary magnesium. For further prospects, see Sec. 4.7.
203
4.7 Conclusions and Outlook The more recent developments using conventional ingot metallurgy (I/M) have resulted in magnesium alloys which are competitive to alternative structural materials. These improvements include: (i) High purity magnesium alloys such as AZ91E with similar or better corrosion resistance compared to most of the commercial aluminum alloys. The AZ alloy series is used for housings in automobiles, aeronautical applications, and terrestrial appliances including computers, mobile telephones, chain saws, etc. (ii) Magnesium alloys with high ductility such as the AM-series are used for energy absorbing applications such as steering wheels, wheels, automotive seat frames, etc. (iii) The AE42 alloy, the AS-series, and the ZC63 alloy show improved elevated temperature properties for engine and transmission components at service temperatures of up to 200 °C. (iv) The new magnesium alloy WE43 exhibits excellent resistance to corrosion and thermal degradation upon longterm exposures to temperatures as high as 300 °C. This alloy is used for aeronautical engine and transmission components, and in racing cars, (v) Superlight Mg-Li-based alloys are particularly aimed as substitutes for computer components and more private appliances usually made of plastics, and where strength, for example, is of secondary concern, (vi) Magnesium-based MMCs for applications at service temperatures as high as 350 °C including engine components such as pistons, for example. Conventional processing has not led to substantial consideration of wrought mag-
204
4 Magnesium-Based Alloys
nesium-based structural products. The scope of rapid solidification processing (RSP) for future wrought magnesium alloys, however, is one of the most exciting areas in materials science, development, and production, and many initiatives are expected to follow. One of the major merits of the more recent work at Allied Signal, within the Pechiney/Norsk collaboration, and at many other places during the last decade was to show that large improvements in strength, ductility, corrosion resistance, and damage tolerance are possible using relatively cheap alloying additions. This includes the use of aluminum and zinc, as well as calcium and strontium, in order to tie-up aluminum in temperature stable, Al 2 Ca dispersions. Thus affording higher aluminum levels at the expense of zinc levels without sacrificing solid solution strengthening in order to arrive at strength levels as for Mg-Al-Zn-based alloys reinforced with refined Mg2Si dispersions, but with a corrosion resistance of the same order of magnitude as that of AZ91E T5, WE43, and EA55RS-type alloys. Despite the progress with conventional and more advanced technologies, magnesium alloys have not yet been considered as a serious alternative by the majority of constructors. The ready availability of RSP magnesium alloys, for example, would necessitate the end user to reconsider the design and use of resultant components for aircraft and automobiles. Drawbacks in the fracture toughness and the available degree of passivity of magnesium engineering alloys have so far prevented them from attracting the attention of the aircraft industry as a potential workhorse for the commercialization of the more recent wrought magnesium alloys. Preliminary insight into the microstructure-property relationship focuses on the universali-
ty of Hall-Petch relationships as the reference for further progress of the achievable properties. Global market forecasts of more concrete possibilities with magnesium are rare, if available at all, leaving many analyses to be carried out. Davis (1991), for example, has undertaken an interesting study on the potential for the substitution of magnesium-based components in established automotive applications (Fig. 4-63). About 12-16% of possible magnesium components were considered to be readily available (Fig. 4-63 b), corresponding to about 17-20% of the potential to save weight with such magnesium-based components (Fig. 4-63 a). Magnesium-based chassis, suspensions, and wheels were considered to cover the largest market potential in automobiles. However, magnesium-based body components were predicted to exceed the potential to save weight via chassis, suspensions, and wheels by 3-4%, resulting in a total of 70 kg (small passenger car) to 120 kg (large car) of magnesium components or 8-10% of an average car's total weight (Fig. 4-63). In Europe, however, only a few new magnesium-based structural applications have yet been recorded. North America has once again taken the lead with an actual growth rate of magnesium alloys of about 7% per year, and forecasts for Asia and particularly for Japan are also very promising (Foley and Gilbert, 1992). With the increasing overall market volume of structural magnesium the volume of dross, sludge, and (oil-) contaminated chips will also increase. Recycling of dross and sludge requires methods to remove harmful impurities such as oil. Techniques under development to separate oil from magnesium include pre-drying via centrifugal methods, subsequent heating to near the boiling point of oil, and evacua-
205
4.7 Conclusions and Outlook M
"60-,
easy to realize
0 9.8 kg
| 15.1 kg
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52.3
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23.0
I
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1.5^ 0 . 0 ^
37.2
J
23.0
S
8.0
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I V*
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30
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Intermediate Car
Small Car
easy to
realize 0 9.8 kg
0 8.7 kg
total
E351 kg
S 50 kg
"c 20-
28.8
14.6 kg
17.8
17.0
II I
14.0 104.8
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13.0
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I
U.3
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rl
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Figure 4-63. (a) Weight of magnesium components and (b) weight reduction potential of using magnesium components in small intermediate, and large cars. A distinction is made between easily realizable substitutions by magnesium and the total potential for substitution (Davis, 1991). [Electr.: electrical & electronic; Transm.: transmission.]
tion of the corresponding oil vapor (Saxena, 1992). Other sources of contamination include coatings for protection of weathered magnesium-based applications against corrosion, and for decoration. New coating developments may allow for more effective recycling procedures, including more thorough elimination of the corresponding coatings and resulting impurities.
Large-scale weight reduction using magnesium alloys requires improved property profiles, for example, while a large-scale weight reduction for automobile applications requires a low and stable price for the alloy and the final product as well. The relative increment in the energy required to produce magnesium compared to that for aluminum and plastics (see Sec. 4.2.1), for example, is largely compensated for by the
206
4 Magnesium-Based Alloys
specific advantages which magnesium offers in terms of alloy transformation, service, and recycling. These include high pressure die casting of magnesium alloys, making near-net-shape production possible during high productivity of complicated components, such as thin-walled door frames, without the need to employ major finishing procedures. An established technology, however, is not necessarily the most effective approach to a forgotten material, while the most readily available alloying approach is not necessarily the most useful approach to a new technology. Also, a new technology is not necessarily the most useful approach to a new alloy transformation principle. Two emerging fields of the more recent developments are spray forming and the formation of metastable magnesium-based alloys offering acceptable fracture toughness, passivation for the first time in history, extremely low corrosion rates, and ultra-high strength values. Research policy for sustainable and more effective development is expected to address these issues more seriously, and to spur some of the more promising developments reviewed by us.
4.8 Acknowledgements The authors are grateful to R. Feser (Metallgesellschaft AG), F. Kruppa (Technical University of Prague), L. Kubin (ONERA-CNRS), U. Mehlhorn (Metallgesellschaft AG), R. Ninomiya (Mitsui Kinzoku), T. Sanders (University of Georgia, Atlanta), and E. E. Schmid (Metallgesellschaft AG) for useful discussions or for their help in preparing this chapter, and to Deutsche Forschungsgemeinschaft for support in preparing Sec. 4.4.
4.9 References Adkins, N.J.E., Tsakiropoulos, P. (1991), Int. J. Rapid Solidif. 6, 87 and 101. Ahmed, M., Lorimer, G. W., Lyon, P., Pilkington R. (1992), in: Magnesium Alloys and Their Applications'. Mordike, B. L., Hehmann, F. (Eds.). Oberursel, F.R.G.: DGM Informationsgesellschaft, pp. 301-308. Aikawa, K., Taketani, K. (1991), E. P. Patent 0407964. Akavipat, S., Hale, E. B., Habermann, C. E., Hagans, P. L. (1985), Mater Sci. Eng. 69, 311. Aluminium Taschenbuch (1984): Aluminium-Zentrale (Ed.). Diisseldorf: Aluminium-Verlag Diisseldorf, p. 56. Ameur, T. B., Yavari, A. R., Barandiaran, J. M. (1991), Mater. Sci. Eng. A134, 1402-1405 Anonymous (1987), Met. Powder Rep. 42, 62. Armstrong, H. H. (1979), The Enigma of the Eighties: Environment, Economics, Energy, Vol. 2. Proc. Natl. SAMPE Symp., San Francisco, CA. Armstrong, J. H., Rawal, S. R., Miscra, M. S. (1990), Mater. Sci. Eng. 126A, 119. ASM (1979), ASM Metals Handbook, Properties of Magnesium Alloys, Vol. 2. Metals Park, OH: ASM Int. ASTM (1984), ASTM StandardB91, Standard Specification for Magnesium Alloy Forging, Annual Book of Standards, Vol.2. Philadelphia, PA: ASTM. Attari, N., Robin, C , Aluvinage, G. (1990), in: Advanced Aluminium and Magnesium Alloys: Khan, T, Effenberg, G. (Eds.). Materials Park, OH: ASM Int., p. 837. Bach, P. (1969), Ph.D. Thesis, Nancy. Badini, C , Marino, F., Montorsi, M., Guo, X. B. (1992), Mater. Sci. Eng. A157, 53-61. Baker, C. F. (1992), in: Magnesium Alloys and Their Applications, Mordike, B. L., Hehmann, F. (Eds.). Oberursel, F.R.G.: DGM Informationsgesellschaft, pp. 77-84. Baliga, C. B. (1991), Ph.D. Thesis, University of Surrey. Baliga, C. B., Tsakiropoulos, P. (1992), in: Magnesium Alloys and Their Applications: Mordike, B. L., Hehmann, F (Eds.). Oberursel, F.R.G.: DGM Informationsgesellschaft, pp. 119-126. Baliga, C. B., Tsakiropoulos, P. (1993), Mater. Sci. Technol. 9, 507 and 513. Baliga, C. B., Tsakiropoulos, P., Watts, J. F , Jeynes, C. (1992), in: Magnesium Alloys and Their Applications: Mordike, B. L., Hehmann, F (Eds.). Oberursel. F.R.G.: DGM Informationsgesellschaft, pp. 191-198. Beck, A. (1940), The Technology of Magnesium and its Alloys. London: Hughes, p. 252. Beer, S., Frommeyer, G., Schmid, E. (1992), in: Magnesium Alloys and Their Applications: Mordike, B. L., Hehmann, F (Eds.). Oberursel, F.R.G.: DGM Informationsgesellschaft, pp. 317-324.
4.9 References
Beer, S., Frommeyer, G., Schmid, E., Helbig, H. (1993), in: Automotive Automation: Croydon, U.K., 26th ISATA Conf., pp. 349-356. Bianchi, L. (1991), J. Met. 5, 45-47. Bode, R., Hornbogen, E. (1992), in: Magnesium Alloys and Their Applications'. Mordike, B. L., Hehmann, F. (Eds.). Oberursel, F.R.G.: DGM Informationsgesellschaft, pp. 461-468. Boettinger, W. I , Bendersky, L., Early, J. G. (1986), Met. Trans. 17A, 781-790. Bollenrath, F. (1939), in: Werkstoff Magnesium. Berlin: VDI-Verlag, p. 10. Bolstad, J. A., Roenhaug, J. B. (1992), in: The Recycling of Metals. Brussels: ASM European Office, pp. 287-302. Borbe, P. C , Erdmann-Jesnitzer, F. (1983), Metall 37, 573. Bothwell, M. R. (1967), in: The Corrosion of Light Metals. Chichester, U.K.: Wiley, p. 84. Bray, D. (1990), AGARD Lecture Series No. 174. Neuilly sur Seine, France: AGARD, paper No. 7. Buhler, K. (1990), Metall 44, 748-753. Busk, R. S. (1952), Trans. AIME 188, 1460. Busk, R. S. (1988), in: Magnesium Products Design. Tokyo: The Japan Magnesium Association, pp. 190-264. Busk, R. S., Leontis, T. E. (1950), Trans. AIME 188, 297-306. Calka, A., Radlinski, A. P. (1989), Mater. Sci. Eng. A118, 131-135. Calka, A., Madhava, M., Polk, D. E., Giessen, B. C. (1977), Scr. Met. 11, 65-70. Carapella, L. A. (1945), Met. Prog. 48, 297-307. Chadwick, G. A., Bloyce, A. (1992), in: Magnesium Alloys and Their Applications: Mordike, B. L., Hehmann, F. (Eds.). Oberursel, F.R.G.: DGM Informationsgesellschaft, pp. 93-100. Chang, C. R, Das, S. K. (1992a), U.S. Patent 5087304. Chang, C. F., Das, S. K. (1992b), U.S. Patent 5129960. Chang, C. F., Das, S. K. (1992c), Private communications, Allied Signal, Morristown, NJ. Chang, C. R, Das, S. K., Raybould, D. (1986), Met. Powder Rep. 41, 302-308. Chang, C. R, Das, S. K., Raybould, D. (1987), in: Processing of Structural Metals by Rapid Solidification: Froes, R H., Savage, S. J. (Eds.). Materals Park, OH: ASM Int., pp. 409-415. Chang, C. R, Das, S. K., Raybould, D., Bye, R. L., Kim, N. J. (1988), in: Space Age Metals Technology: Froes, F . H , Cull. R. A. (Eds.). Covina, CA: Society for the Advancement of Materials and Process Engineering (SAMPE), pp. 345-356. Chang, C. R, Das, S. K., Raybould, D., Bye, R. L., Limoncelli, E. V. (1989), in: 1989 Advances in Powder Metallurgy: Princeton, NJ: Metal Powder Industries Federation. Chang, C. R, Das, S. K., Raybould, D., Limoncelli, E. V. (1991), in: PM in Aerospace and Defence Tech-
207
nologies: Froes, R H. (Ed.). Princeton, NJ: Metal Powder Industries Federation (MPIF), pp. 393400. Chen, N. K., Pond, R. B. (1952), Trans. AIME 194, 1085. Chin, E. S. C , Nunes, J. (1988), in: Advances in Magnesium Alloys and Composites: Paris, P. G., Hunt, W. H. (Eds.). Warrendale, PA: TMS, p. 123. Chisholm, D. S., Busk, R. S. (1953), U.S. Patent 2630 623. Chisholm, D. S., Hershey, G. R (1956), U.S. Patent 2752196. Clark, J. B. (1962), Ada Met. 16, 141. Clark, J. B. (1986), Trans. JIM 9, 354. Clyne, T. W, Ricks, R. A., Goodhew, P. J. (1984), Int. J. Rapid Solidif 1, 59-80. Colby, N. R., Hershey, G. R (1955), U.S. Patent 2699 576. Cotton, J. D., Jones, H. (1989), /. Electrochem. Soc. 136 (11), 523C-527C. Cotton, J. D., Jones, H. (1991), Int. J. Rapid Solidif. 6, 155. Cottrell, A. (1975), An Introduction to Metallurgy, 2nd ed. Sevenoaks, U.K.: Edward Arnold, p. 398. Couling, S. L: (1985), private communication, Batelle Laboratories, Columbus, OH. Courtney, T. H. (1990), Mechanical Behaviour of Materials. New York: McGraw-Hill. Cutler, H. (1992), in: Magnesium on the Move, 49th World Magnesium Conference. McLean, VA: International Magnesium Association, pp. 71-78. Das, S. K. (1993), private communication. Das, S. K., Chang, C P. (1989), U.S. Patent 4853035. Das, S. K., Davis, L. A. (1988), Mater. Sci. Eng. 98, 8-12. Das, S. K., Raybould, D., Bye, R. L., Chang, C. R (1988), U.S. Patent 4718475. Das, S. K., Chang, C. R, Raybould, D. (1989), in: PM in Aerospace and Defense Technologies', Froes, R H. (Ed.). Princeton, NJ: Metal Powder Industries Pederation (MPIF), pp. 63-66. Das, S. K., Chang, C. R, Raybould, D. (1990), U.S. Patent 4938 809. Das, S. K., Chang, C. R, Raybould, D. (1991), Light Met. Age, June, 41. Das, S. K., Chang, C. R, Raybould, D., King, J. P., Thistlethwaite, S. (1992), in: Magnesium Alloys and Their Applications: Mordike, B. L., Hehmann, P. (Eds.). Oberursel, F.R.G.: DGM Informationsgesellschaft, pp. 487-494. Davis, J. (1991), SAE Technical Paper Series, No. 910551. Dean, R. S., Anderson, C. T. (1943), U.S. Patent 2 317 980. Dieter, G. E. (1986), Mechanical Metallurgy, 3rd ed. Singapore: McGraw-Hill, p. 137. Diwanji, A. P., Hall, I. W. (1987), in: Proc. American Society for Composites, 2nd Tech. Mtg. Lancaster, PA: Technomic, pp. 351-360.
208
4 Magnesium-Based Alloys
Dow Chemical Company (1962 a), Low Temperature Properties of Wrought Magnesium Alloys, TS&D Letter Enclosure, March 23. Dow Chemical Company (1962 b), Magnesium Forging Alloys for Elevated Temperature Service, TS&D Letter Enclosure, April 24. Dow Chemical Company (1964a), HM12A-T8 Magnesium Alloy Sheet & Plate, TS&D Letter Enclosure, April 24 Dow Chemical Company (1964 b), ZE10A Sheet & Plate Magnesium Alloy, TS&D Letter Enclosure, June 1. Dow Chemical Company (1964 c), HK31A-H24 and -O Sheet & Plate Magnesium Alloy, TS&D Letter Enclosure, July 6. Dow Chemical Company (1965), Magnesium Mill Products. Dow Chemical Company (1967), Magnesium in Design, Bulletin No. 141-213. Dow Chemical Company (1983 a), Magnesium Mill Products and Alloys, Bulletin No. 141-481. Dow Chemical Company (1983 b), Magnesium Tooling Plate, Bulletin No. 141-232. Dries, G. A., Tompkins, S. S. (1988), in: Metal Matrix, Carbon and Ceramic Matrix Composites (Conf Proc), Cocoa Beach, FL, NASA CP-3018, p. 97. Drits, M. E., Sviderkaya, Z. A., Rokhlin, L. L., Nikitina, N. I. (1979), Metalloved. Term. Obrab. Met. 11, 62. Drits, M. E., Rokhlin, L. L., Abrukina, N. P. (1985), Metalloved. Term. Obrab. Met. 17, 27-28. Elias, L. G., Duszcyk, X, Hehmann, F. (1992), in: Magnesium Alloys and Their Applications: Mordike, B. L., Hehmann, F. (Eds.). Oberursel, F.R.G.: DGM Informationsgesellschaft, pp. 343-350. Emley, E. F. (1966), Principles of Magnesium Technology. Oxford: Pergamon. Faure, J. F , Nussbaum, G., Regazzoni, G. (1991), Eur. Patent 414620 Al, U.S. Patent 5073 207. Fenn, R. W. Jr. (1961), in: Symp. on Low Temperature Properties of High-Strength Aircraft and Missile Materials, STP No. 287: Metals Park, OH: ASTM. Fenn, R. W. Jr., Lockwood, L. F. (1960), Welding J. {Miami), 352. Fluck, E., Heumann, K. G. (1986), Periodic Table of the Elements. Weinheim, F.R.G.: VCH. Foerster, G. S., Johnson, H. A. (1958), Prod. Eng. 29 (19), 80-81. Foley, P. T, Gilbert, K. (1992), in: Magnesium on the Move, 49th World Magnesium Conference. McLean, VA: International Magnesium Association, pp. 7-16. Fox, F. A., Lardner, E. (1943), J. Inst. Met. 49, 373. Froats, A., Aune, T. K., Hawke, D., Unsworth, W, Hillis, J: (1987), in: ASM Metals Handbook, Vol. 13, 9th ed.: Metals Park, OH: ASM Int. pp. 740754. Froes, F. H., Kim, Y. W, Hehmann, F. (1987). J. Met. 39(8), 14-21.
Froes, F. H., Quist, W. E., Das, S. K. (1988), Met. Powder Rep. 43, 401-404. Garber, F. (1993), Metall 47, 574-577. Garboggini, A., McShane, H. B. (1992), in: Magnesium Alloys and Their Applications: Mordike, B. L., Hehmann, F (Eds.). Oberursel, F.R.G.: DGM Informationsgesellschaft, pp. 503-510. Gjestland, H., Nussbaum, G., Regazzoni, G. (1989), in: Light Weight Alloys for Aerospace Applications: Lee, E.W., Chia, E. H., Kim, N. J. (Eds.). Warrendale, PA: TMS, pp. 139-149. Gjestland, H., Nussbaum, G., Regazzoni, G., Lohne, O., Bauger, O. (1991), Mater. Sci. Eng. A134, 1197-1200. Goddard, D. M., Whitman, W. R., Pumphrey, R. L. (1986), SAE Technical Paper Series 860564. Gonzalez-Doncel, G., Wolfenstine, X, Metenier, P., Ruano, O. A., Sherby, O. D. (1990), /. Mater. Sci., 25, 4535. Graebner, X E., Chen, H. S. (1987), Phys. Rev. Lett. 58, 1945. Grensing, F. C , Fraser, H. L. (1987), in: Processing of Structural Metals by Rapid Solidification: Froes, F. H., Savage, S. X (Eds.). Metals Park, OH: AMS Int., pp. 429-435. Grube, G., von Zeppelin, H., Bumm, H. (1934), Z. Elektrochem. 40, 160-164. Gwynne, B., King, X F., Unsworth, W (1988), in: Magnesium Developments: International Magnesium Association: McLean, VA, p. 1. Haferkamp, H., Bach, F.-W, Bohling, P., Willems, C. (1990), in: Advanced Aluminium and Magnesium Alloys: Khan, T, Effenberg, G. (Eds.). Metals Park, OH: ASM Int., pp.829-836. Haferkamp, H., Bach, F.-W, Bohling, P., Juchmann, P. (1992), in: Magnesium Alloys and Their Applications: Mordike, B. L., Hehmann, E. (Eds.). Oberursel, F.R.G.: DGM Informationsgesellschaft, pp. 243-250. Haferkamp, H., Bohling, P., Juchmann, P. (1993), Forschung, Mitteilungen der DFG 2/93, 24. Hafner, X (1976), / Phys. K, Met. Phys. 6, 1243. Hall, E. O. (1970), in: Yield Point Phenomena in Metals and Alloys. New York: Plenum. Hall, I. W (1987), Scr. Metall. 21, 1717. Hall, I. W (1991), J. Mater. Sci. 26, 776. Hanawalt, X D., Nelson, C. E., Peloubet, X A. (1942), Trans. AIME 147, 273-299. Hansen, F., Schiirmann, E., Frommeyer, G. (1986), Metall 40, 146-148. Haughton, X L., Prytherch, W E. (1937), Magnesium and its Alloys, London: HMSO. Hauser, F. E., Landon, P. R., Dora, X E. (1958), Trans. ASM 50, 856. Hazelton, L. E. (1991), U.S. Patent 5074936. Hehmann, F. (1984), M. Met. Thesis, University of Sheffield. Hehmann, F. (1985), Final Report on SUMAC Job No. 37401, University of Sheffield.
4.9 References
Hehmann, F. (1986), Final Report on Contract F 33615-84-C-5127, Task Nos. 28 and 54, WrightPatterson AFB. Hehmann, F. (1987), Final Report on 3-MonthProject at The University of Sheffield, MaxPlanck-Institut, Stuttgart. Hehmann, F. (1988), Doctoral Thesis, The University of Stuttgart. Hehmann, F. (1990a), Acta Metall. Mater. 38, 979992. Hehmann, F. (1990 b), in: Proc. of the 47th International Magnesium Conference, May 29-31, Cannes. McLean, VA: Int. Magnesium Association, pp. 76-82. Hehmann, F. (1992 a), in: Magnesium Alloys and Their Applications: Mordike, B. L., Hehmann, F. (Eds.). Oberursel, F.R.G.: DGM Informationsgesellschaft, pp. 21-28. Hehmann, F. (1992 b), in: Conf Proc. on Advances in Production and Fabrication of Light Metals and Metal Matrix Composites: Avedesian, M. M. (Ed.). Montreal: Canadian Institute of Mining, Metallurgy and Petroleum (CIMMP), pp. 305-320. Hehmann, F. (1995 a), Work in progress. Hehmann, F. (1995 b), Mater. Sci. Rev., in preparation. Hehmann, F , Jones, H. (1986), in: Rapidly Solidified Alloys and Their Mechanical and Magnetic Properties: Pittsburgh, PA: Mater. Res. Soc, pp. 259-274. Hehmann, F , Jones, H. (1987), in: Magnesium Technology: London: The Inst. of Metals, pp. 83-96. Hehmann, F , Jones, H. (1990), in: Advanced Aluminium and Magnesium Alloys: Khan, T, Effenberg, G. (Eds.). Metals Park, OH: ASM Int., 781-788. Hehmann, F , Jones, H. (1993), in: Rapid Solidification Technology - An Engineering Guide: Srivatsan, T. S., Sudarshan, T. S. (Eds.). Lancaster, PA: Technomic, pp. 441-487. Hehmann, F , Tsakiropoulos, P. (1992), in: Magnesium Alloys and Their Applications', Mordike, B. L., Hehmann, F. (Eds.). Oberursel, F.R.G.: DGM Informationsgesellschaft, pp. 535-542. Hehmann, F , Krishnamurthy, S., Robertson, E., Jackson, A. G., Savage, S. J., Froes, F. H. (1986), in: Horizons of Powder Metallurgy, Part II: Freiburg: Verlag Schmid, pp. 1001-1008. Hehmann, F , Krishnamurthy, S., Robertson, E., Savage, S. J., Froes, F. H., USAF (1988), U.S. Patent 4770 850. Hehmann, F., Sommer, F , Jones, H., Edyvean, R. G. J. (1989), /. Mater. Sci. 24, 2369-2379. Hehmann, F , Sommer, F , Predel, B. (1990), Mater. Sci. Eng. A125 (2), 249-265. Higashi, K., Wolfenstine, J. (1991), Mater. Lett. 10, 329. Higashi, K., Kubota, K., Neite, G. (1992), in: Magnesium Alloys and Their Applications: Mordike, B. L., Hehmann, F. (Eds.). Oberursel, F.R.G.: DGM Informationsgesellschaft, p. 293.
209
Higashi, K., Aoki, H., Kubota, K., Tanimura, S. (1993), in: Aspects of High Temperature Deformation and Fracture in Crystalline Materials: Hosoi, Y, Yoshinaga, H., Oikawa, H., Maruyama, K., (Eds.). Sendai: The Japan Institute of Metals, p. 479. Hino, H., Komatsu, M., Mori, H. (1990), / Jpn. Inst. Light Met. 40, 123. Hirota, E., Habazaki, H., Kawashima, A., Asami, K., Hasimoto, K. (1993), Scientific Report A38, The University of Tohuko, Japan. Hirth, J. P., Lothe, J. (1968), Theory of Dislocations. New York: McGraw-Hill. Hollrigl-Rosta, F , Just, E., Kohler, J., Melzer, H. J. (1980), in: Proc. 37th World Magnesium Conference. McLean, VA: International Magnesium Association, pp. 38-45. Hopkins, W. G. (1991-1993), Private communications, PSI Ltd. Research & Pilot Plant Systems, Polegate, East Sussex. Housh, S. E., Petrovich, V. (1992), SAE Technical Paper Series 920071. Housh, S. E., Mikucki, B., Stevenson, A. (1990), in: ASM Metals Handbook, Vol. 2, 10th ed.: Metals Park, OH: ASM Int., pp. 455-479. Inoue, A., Nakano, U., Bizen, Y, Masumoto, T., Chen, H. S. (1988 a), Jpn. J. AppL Phys. 27 (6), L944-L947. Inoue, A., Ohtera, K., Kita, K., Masumoto, T. (1988 b), Jpn. J AppL Phys. 27(12), L2248-L2251. Inoue, A., Kohinata, M., Tsai, A. P., Masumoto, T. (1989), Mater. Trans. JIM 30 (5), 378-381. Inoue, A., Kato, A., Zhang, T, Kim, S. G., Masumoto, T. (1991), Mater. Trans. JIM 32 (7), 609-616. International Magnesium Association (1993), personal communication by B. Clow. Isserow, S., Rizzitano, J: (1974), Int. J. Powder Metall. Powder Technol. 109 (3), 217. Ivanov, E., Bokhonov, B., Konstanchuk, I. (1987), J. Mater. Sci. 26, 1409. Jackson, J. H., Frost, P. D., Loonam, A. C , Eastwood, L. W, Lorig, C. H. (1949), /. Metals, Trans. AIME 185, 149-168. Jackson, R. X, Frost, P. D. (1967), NASA Report SP-5068. Johnston, W. G., Gilman, J. J. (1957), J. AppL Phys. 30, 121. Jones, H. (1989), Int. J. Rapid Solidif 4, 297-315. Jones, H. (1991), Mater. Sci. Eng. A137, 77-85. Jones, H. (1993), in: Encyclopedia of Materials Science Engineering, Suppl. Vol. 3: Cahn, R. W. (Ed.). Oxford: Pergamon. Jones, H., Joshi, A., Rowe, R. G., Froes, F. H. (1987), Int. J Powder Metall. 23, 13-24. Joshi, A., Lewis, R. E. (1988), in: Advances in Magnesium Alloys: Paris, H., Hunt, W. H. (Eds.). Warrendale, PA: TMS, pp. 89-103. Joshi, A., Lewis, R. E., Jones, H. (1989), Int. J. Rapid Solidif 4, 251-269.
210
4 Magnesium-Based Alloys
Joshi, A., Adamson, R. D., Lewis, R. E. (1992), in: Magnesium Alloys and Their Applications: Mordike,B. L.,Hehmann, F. (Eds). Oberursel, F.R.G.: DGM Informationsgesellschaft, pp. 495-502. Juarez-Islad, J. A. (1991), Mater. Sci. Eng. A134, 1193-1196. Kainer, K. U. (1989), in: Proc. Advanced Materials and Processes - EUROMAT '89: Exner, H. E., Schumacher, V. (Eds.). Vol. 2, Aachen, 22-24 Nov. Oberursel, F.R.G.: DGM Informationsgesellschaft, pp. 1331-1336. Kainer, K. U. (1990), Met. Powder Rep. 45, 684-687. Kainer, K. U. (1991 a), in: Advances in Powder Metallurgy - 1991, Vol. 6, Chicago, IL, June 9-12. Princeton, NJ: Metal Powder Industries Federation, pp. 307-321. Kainer, K. U. (1991b), in: Metal Matrix Composites - Processing, Microstructure and Properties'. BildeSorensen, J. B., Hansen, N., Horsewell, A., Leffers, T, Lilholt, H. (Eds.). Roskilde, Denmark: Ris0 National Laboratory, p. 429. Kainer, K. U. (1992), in: Magnesium Alloys and Their Applications: Mordike, B. L., Hehmann, F. (Eds.). Oberursel, F.R.G.: DGM Informationsgesellschaft, pp. 415-422. Kainer, K. U., Schroder, J. (1990), in: Int. Conf. on Powder Metallurgy: London: Inst. of Metals, 398. Kainer, K. U., Schroder, X, Mordike, B. L. (1991), in: PM Aerospace Materials 1991, Lausanne, Switzerland, Nov. 4-6: Shrewsbury, U.K.: MPR Publishing Services, paper No. 38. Kalimullin, R. K. H., Berdnikov, A. T. (1986), Zashch. Met. 22, 262-264. Kalimullin, R. K. H., Valuev, V. V, Bernikov, A. T. V. (1986), Metalloved. Term. Obrab. Met. 9, 39-40. Kalimullin, R. K. H., Spiridonov, V. B., Bernikov, A. T., Romanov, A. A., Pantikina, G. N. (1988), Metalloved-Term. Obrab. Met. 5, 18-24. Karney, R., Sachs, G. (1928), Z. Phys. 49 480. Kato, K., Tokisue, H., Chiba, I. (1982), /. Jpn. Inst. Light Met. 42 (8), 453-458. Kieschke, R. R., Warwick, C. M., Clyne, T. W. (1991), Acta Metall. Mater. 39, 445-452. Kindler, H., Nikles, A. (1980), Kunststoffe 70, 802807. King, J. F , Hehmann, F. (1992), in: Magnesium Alloys and Their Applications: Mordike, B. L., Hehmann, F (Eds.). Oberursel, F.R.G.: DGM Informationsgesellschaft, pp. 309-316. King, J. F , Wardlow, G. D. (1993), private communication. Kirk-Othmer (1981), Encyclopedia of Chemical Technology, Vol. 14, 3rd ed. New York: Wiley. Kittel, C. (1976), Introduction to Solid State Physics, 5th ed. New York: Wiley. Kittilsen, B., Pinfold, M. D. (1992), in: Magnesium Alloys and Their Applications: Mordike, B. L., Hehmann, F. (Eds.). Oberursel, F.R.G.: DGM Informationsgesellschaft, pp. 85-92.
Kleine, A., Dudek, H. J. (1992), in: Magnesium Alloys and Their Applications: Mordike, B. L., Hehmann, F. (Eds.). Oberursel, F.R.G.: DGM Informationsgesellschaft, pp. 447-454. Krishnamurthy, S., Kim, Y W. (1988), in: Magnesium Developments, Proc. World Materials Congress, Sept., Chicago IMA. Metals Park, OH: ASM Int., pp. 11-16. Krishnamurthy, S., Khobaib, M., Robertson, E., Froes, F. H. (1988), Mater. Sci. Eng. 99, 507-511. Krishnamurthy, S., Weiss, I., Froes, F H. (1989), Key Eng. Mater. 29-31, 135-146. Kurz, W, Giovanola, B., Trivedi, R. (1987), Acta Metall. 35, 823. Lahaie, D., Embury, J. D., Chadwick, M. M., Giray, G. T. (1992), Scr. Metall. Mater. 27, 139-142. Landolt-Bornstein (1979), New Series, Group III: Crystal and Solid State Physics, Vol. 1. Berlin: Springer. Laurent, V, Jarry, P., Regazzoni, G., Apelian, D. (1992), J. Mater. Sci. 27, 4447. Lavernia, E., Gomez, E., Grant, N. J. (1987), Mater. Sci. Eng. 95, 225-236. Lavernia, E. J., Ayers, J. D., Srivatsan, T. S. (1993), in: Rapid Solidification Technology - An Engineering Guide: Lancaster, PA: Technomic, p. 315. Lee, R. E., Jones, W. J. D. (1974), J. Mater. Sci. 9, 469-475. Leontis, T. E. (1949), Trans. AIME 185, 968. Leontis, T. E., Busk, R. S. (1953), G.B. Patent 690783. Leontis, T. E., Nelson, C. E. (1951), Trans. AIME 191, 120. Lewis, R. E., Joshi, A. (1989), in: PM in Aerospace and Defense Technologies: Froes, F H. (Ed.). Princeton, NJ: Metal Powder Industries Federation (MPIF), pp. 55-62. Lewis, R. E., Joshi, A., Jones, H. (1987), in: Processing of Structural Metals by Rapid Solidification: Froes, F H., Savage, S. J. (Eds.). Metals Park, OH: AMS Int., pp. 367-378. Li, Y, Jones, H. (1992), in: Magnesium Alloys and Their Applications: Mordike, B. L., Hehmann, F. (Eds.). Oberursel, F.R.G.: DGM Informationsgesellschaft, pp. 543-550. Li, Y, Liu, H. Y, Jones, H. (1992a), in: Magnesium Alloys and Their Applications: Mordike, B. L., Hehmann, F (Eds.). Oberursel, F.R.G.: DGM Informationsgesellschaft, pp. 479-486. Li, Y, Davies, H. A., Jones, H. (1992b), in: Magnesium Alloys and Their Applications: Mordike, B. L., Hehmann, F. (Eds.). Oberursel, F.R.G.: DGM Informationsgesellschaft, pp. 551-557. Lim, S., Choh, T. (1991), in: Int. Conf. on Recent Advances in Science and Engineering of Light Metals, Sendai, Japan, p. 175. Lipton, P., Kurz, W, Trivedi R. (1987), Acta Metall. 35, 957.
4.9 References
Lohne, O., Bauger, O., Gjestland, H., Nussbaum, G., Regazzoni, G. (1991), in: Science and Engineering of Light Metals, Proc. Int. Conf. RASELM, JILM, Tokyo: pp. 163-168. Magnesium Die Casting Manual (1993), Dow Chemical (Japanese). Magnesium Elektron Ltd. (1961), ZTY, A Wrought Magnesium Alloy for Use at Elevated Temperature. Magnesium Handbook (1975). Tokyo: Japan Magnesium Association, pp. 139-160. Magnesium Manual (1991). Tokyo: Japan Magnesium Association, p. 42. Magnesium Manual (1992). Tokyo: Japan Magnesium Association, p. 90. Makar, G. L. (1991), Doctoral Dissertation, Johns Hopkins University. Makar, G. L., Kriiger, J. (1993), Int. Mater. Rev. 38, 138-153. Makar, G. L., Kriiger, J., Joshi, A. (1988), in: Advances in Magnesium Alloys: Paris, H., Hunt, W. H. (Eds.). Warrendale, PA: TMS, pp. 105-121. Makar, G. L., Kriiger, I , Sieradski, U. (1992), /. Electrochem. Soc. 139 (1), 47-53. Mason, J. R, Warwick, C. M., Smith, P. J., Charles, J. A., Clyne, T. W. (1989), /. Mater. Sci. 24, 3934. Massalski, T. B., Okamoto, H., Subramanian, P. R., Kacprzak, L. (Eds.) (1990), Binary Alloy Phase Diagrams. Metals Park, OH: ASM Int. Masumoto, T, Inoue, A., Sakuma, T., Shibata, T. (1992), U.S. Patent 5118 368. McDonald, J. C. (1968), Trans. ASM 61, 505. Mercer, W. E., II (1990), SAE Technical Paper Series 900788. Meschter, P. J. (1986), Private communication, McDonell Douglas Research Laboratories. Meschter, P. J. (1987), Met. Trans. 18A, 347-350. Meschter, P. J., O'Neal, J. E. (1984), Met. Trans. 15 A, 234-240. Metallstatistik (1990), Metallgesellschaft AG, Frankfurt am Main, F.R.G. Metenier, P., Gonzalez-Doncel, G., Ruano, O. A., Wolfenstine, J., Sherby, O. D. (1990), Mater. Sci. Eng.A125, 195. Mikucki, B. A., Mercer, W. E. II, Green, W. G. (1990), SAE Technical Paper Series 900533. Muller, C. (1991), Metall 45, 934-936. Miiller, M., Wachter, I , Sommer, F. (1992), in: Magnesium Alloys and Their Applications: Mordike, B. L., Hehmann, F. (Eds.). Oberursel, F.R.G.: DGM Informationsgesellschaft, pp. 527-534. Murray, R. W., Hillis, J. E. (1990), SAE Technical Paper Series, No. 900791. Nabarro, F. R. N. (1967), Theory of Crystal Dislocations. Oxford: Clarendon. Neite, G., Rudolph, G., Schmid, E. E., Schulze, T, Kubota, K., Ninomiya, R. (1993), Published Application WO 93/19216. Ninomiya, R., Kubota, K., Ojiro, T. (1995), Acta Metall. Mater. 43, 669.
211
Nussbaum, G., Gjestland, H., Regazzoni, G. (1989 a), in: Proc. of the 45th Annual World Magnesium Conference, Dearborn, MI. McLean, VA: Int. Magnesium Assoc, pp. 72-78. Nussbaum, G., Saintfort, P., Regazzoni G., Gjestland, H. (1989b), Scr. Met. 23, 1079. Nussbaum, G., Regazzoni, G., Gjestland, H. (1990 a), in: Proc. of the 47th International Magnesium Conference, May 29-31, Cannes. McLean, VA: Int. Magnesium Assoc, pp. 130-135. Nussbaum, G., Regazzoni, G., Gjestland, H. (1990 b), in: Advanced Aluminium and Magnesium Alloys: Khan, T, Effenberg, G. (Eds.). Metals Park, OH: ASM Int., pp. 789-796. Nussbaum, G., Reggazoni, G., Gjestland, H. (1991), in: Science and Engineering of Light Metals, Proc. Int. Conf RASELM, JILM, Tokyo: pp. 114-120. Olsen, A. LO. (1992), Metall 46, 570-574. 0ymo, D., Holta, O., Hustoft, O.-M., Henricksson, J. (1992), Metall 46, 898. Pan, L. M., Saunders, N., Tsakiropoulos, P. (1989), Mater. Sci. Technol. 5, 609-612. Pane, R. J. M., Bailey, N. (1959), J. Inst. Met. 88,417. Pekgiileryiiz, M. O. (1991), in: Proc. of the 1st Canadian Int. Composite Conf, Montreal, Canada. Pekgiileryiiz, M. O., Avedesian, M. M. (1992), in: Magnesium Alloys and Their Applications; Mordike, B. L., Hehmann, F. (Eds.). Oberursel, F.R.G.: DGM Informationsgesellschaft, pp. 213-220. Pekgiileryiiz, M. O., Luo, A., Vermetta, P., Avedesian, M. M. (1993), in: Proc. of the 50th Int. Magnesium Conf. McLean, VA: Int. Magnesium Assoc. Pennington, J. N. (1993), Mod. Met. 49, Sept., 34L34P. Pickering, H. W, Wagner, C. (1967), /. Electrochem. Soc. 114, 698. Polmear, I. J. (1992), in: Magnesium Alloys and Their Applications: Mordike, B. L., Hehmann, F. (Eds.). Oberursel, F.R.G.: DGM Informationsgesellschaft, pp. 201-212. Pourbaix, M. (1974), Atlas of Electrochemical Equilibria in Aqueous Solutions. Houston, TX: National Association of Corrosion Engineers, p. 141. Project MAIE/0053/(1990), F: Rapid Solidification of Magnesium, EURAM-BRITE Workshop, 26-27 Nov., Louvain-la-Neuve, Belgium. Rabinovitch, M., Daux, J. C , Raviort, J. L., VidalSetif, M. H., Mevrel, R., Abiven, H., Peltier, J. F. (1992), in: Proc. Int. Symp. on Advanced Materials for Lightweight Structures, ESTEC, Noordwijk, The Netherlands. Raybould, D., Chang, C. F., Das, S. K. (1991), U.S. Patent 5071474. Raynor, G. V. (1959), The Physical Metallurgy of Magnesium and Its Alloys. London: Pergamon. Regazzoni, G., Nussbaum, G., Gjestland, H. (1991), U. S. Patent 4997622. Reichek, K. N., Clark, K. I , Hillis, J. E. (1985), SAE Technical Paper Series, No. 850417.
212
4 Magnesium-Based Alloys
Roberts, C. S. (1960), Magnesium and Its Alloys. New York: Wiley. Robinson, H. A., George, P. F. (1954), Corrosion 10, 182-188. Rohklin, L. L., Oreshkina, A. A., Dutkevich, Y, Sotsyush-Podosek, M. (1990), Russ. Metall (Engl. Transl.)4, 116-119. Rohklin, L. L., Dobatkina, T. V., Korol'kova, I. G., Grin, Y N. (1991), Russ. Metall. {Engl. Transl.) 5, 182-185. Sattler, H.-P., Yoshida, T. (1992), in: Magnesium Alloys and Their Applications: Mordike, B. L., Hehmann, F. (Eds.). Oberursel, F.R.G.: DGM Informationsgesellschaft, pp. 29-36. Sauerwald, F. (1947), Anorg. Z. Allg. Chem. 225, 212. Saunders, N., Tsakiropoulos, P. (1988), Mater. Sci. Technol. 4, 157-160. Saxena, S. K. (1992), Br. Patent 2257162. Sayashi, M., Hino, H., Komatsu, M., Sasaki, M: (1988), in: How to Apply Advanced Composite Technology, Conf. Proc, Dearborn, MI. Metals Park, OH: ASM International, p. 479. Schemme, K. (1993 a), Doctoral Thesis, University of Bochum, Diisseldorf: VDI-Verlag. Schemme, K. (1993 b), Private communication. Schemme, K., Hornbogen, E. (1990), in: Advanced Aluminium and Magnesium Alloys: Khan, T., Effenberg, G. (Eds.). Metals Park: OH: ASM Int., pp. 797-803. Schemme, K., Hornbogen, E. (1992), in: Magnesium Alloys and Their Applications; Mordike, B. L., Hehmann, F. (Eds.). Oberursel, F.R.G.: DGM Informationsgesellschaft, pp. 519-525. Schroder, X, Kainer, K. U. (1991), Mater. Sci. Eng. A135, 33. Schurmann, E., Engel, R. (1986), Giessereiforschung 38, 58-72. Schurmann, E., Hansen, F. (1982), Metall 36, 645648. Serve, G. (1962), VDI-Report No. 58, p. 71. Shamrai, F. I. (1947), Bull. Acad. Sci., U.S.S.R., Sec. Chem. Sci. 605-616; (1948) 83-94 and 290-301. Shao, G., Tsakiropoulos, P., Mivdownik, A. P. (1993), Int. J. Rapid Solidif. 8, 41-68. Smithells, C. J. (1976), Metals Reference Book. London: Butterworths. Solberg, J. K., Torklep, J., Bauger, O., Gjestland, H. (1991), Mater. Sci. Eng. A134, 1201-1203. Sommer, F. (1993), private communication. Sommer, F., Fripan, M., Predel, B. (1982a), in: 4th Int. Conf. on Rapidly Quenched Metals, Sendai: Masumoto, T, Suzuki, K. (Eds.). Tokyo: The Japan Inst. Metals, pp. 209-212. Sommer, F., Vogelbein, W, Predel, B. (1982b), /. Non-Cryst. Solids 51, 333-343. Sommer, R, Hehmann, F , Jones, H. (1990), /. LessCommon Met. 159, 237-259. Stohr, J. F., Poirier, J. P. (1972), Phil. Mag. 25,1313. Stall, D. R., Sinke, G. C. (1956), Thermodynamic Properties of the Elements, Advances in Chemistry
Series 18. Washington, DC: American Chemical Society, p. 124. Subramanian, R., Sircar, S., Mazumdar, J. (1991), /. Mater. Sci. 26, 951-956. Taleff, E. M., Ruano, O. A., Wolfenstine, J., Sherby, O. D. (1992), J. Mater. Res. 7, 2131. Tammann, G. (1919), Z. Anorg. Allg. Chem. 107, 1-237. Thorbeck, X, Hehmann, F. (1996), unpublished. Tischer, R. P., Gerischer, H. (1958), Z. Electrochem. 62, 50. Touloukian, Y S. Kirby, R. K., Taylor, R. E., Desai, P. D. (1978), Thermophysical Properties of Matter. New York: IFI/Plenum. Unal, A. (1992), Mater. Manuf. Processes 7,441-461. Unsworth, W (1989), Int. J. Mater. Product Technol. 4, 359. Vervoort, P. X, Duszczyk, X (1991), in: PM Aerospace Materials, Lausanne, Switzerland, Nov. 4-6: Shrewsbury, U.K.: MPR Publishing Services, paper No. 30. Viala, J. C , Fortier, P., Clareyrolas, G., Vincent, H., Bouix, X (1991), J. Mater. Sci. 26, 4977. Vitek, V., Igarshi, M. (1991), Phil. Mag. 63 A, 10591075. VKE (1993), Private communication, Verband der Kunststofferzeugenden Industrie e.V., Frankfurt, von Mises, R. (1928), Z. Angew. Math. Mech. 8, 161. Vostry, P., Stulikova, I., Smola, B., Riehemann, W, Mordike, B. L. (1991), Met. Sci. Eng. A137, 8792. Vostry, P., Stulikova, I., Smola, B., Riehemann, W, Mordike, B. L. (1992), in: Magnesium Alloys and Their Applications: Mordike B. L., Hehmann, F. (Eds.). Oberursel, F.R.G.: DGM Informationsgesellschaft, pp. 511-517. Whitby, L. (1963), in: Corrosion Resistance of Metals and Alloys: McKay, R. X, LaQue, F. L. (Eds.). New York: Reinhold Publishing, p. 169. Wilks, T. E., King, X F. (1992), in: Magnesium Alloys and Their Applications: Mordike, B. L., Hehmann, F. (Eds.). Oberursel, F.R.G.: DGM Informationsgesellschaft, pp. 431-437. Zelin, M. G., Valiev, R. Z., Yang, H. S., Mukherjee, A. K. (1991), Ser. Metall. Mater. 25, 2775-2780.
General Reading Bush, R. S. (1987), Magnesium Product Design. New York: Marcel Dekker. Emley, E. F. (1966), Principles of Magnesium Technology. Oxford: Pergamon. Kurz, W, Fisher, D. X (1989), Fundamentals of Solidification. Aedermannsdorf, Switzerland: Trans Tech. Raynor, G. V. (1959), The Physical Metallurgy of Magnesium and Its Alloys. Oxford: Pergamon. Srivatsan, T. S., Sudarshan, T. S. (1993), Rapid Solidification Technology. Lancaster, PA: Technomic.
5 Aluminum-Based Alloys Yotaro Murakami The New Materials Center, Osaka Science and Technology Center, Osaka, Japan
List of 5.1 5.2 5.2.1 5.2.2 5.2.2.1 5.2.2.2 5.2.2.3 5.2.3 5.2.3.1 5.2.3.2 5.2.3.3 5.2.3.4 5.2.3.5 5.2.4 5.2.4.1 5.2.4.2 5.2.4.3 5.2.5 5.2.5.1 5.2.5.2 5.2.5.3 5.3 5.3.1 5.3.2 5.3.3 5.3.4 5.3.5 5.3.5.1 5.3.5.2 5.4 5.4.1 5.4.1.1 5.4.1.2 5.4.2 5.4.3
Symbols and Abbreviations Introduction Intrinsic Characteristics of Aluminum Lightness Tensile Strength Solid-Solution Hardening Work Hardening Precipitation Hardening Precipitate Structures and Technological Properties Bimodal Precipitate Structures and Grain Refinement Precipitate Structures and Fracture Toughness Precipitate Structures and Fatigue Strength Precipitate Structures and Fine-Grain Superplasticity Precipitate Structures and Stress-Corrosion Cracking Physical Properties Relating to Commercial Applications Electrical Resistivity Thermal Conductivity Other Physical Properties Corrosion Characteristics Surface Oxide Film Contact Corrosion Behavior with Other Metals Influence of Alloying Elements on Corrosion Advantages in Manufacturing and Fabrication Processes Melting and Casting Operations Versatile Deformation Techniques Good Machinability Versatile Bonding Technologies Advanced Processing Technologies Rapid Solidification Processes (RSPs) Mechanical Alloying Process (MAP) Main Aluminum Alloy Groups Alloy and Temper Designations Alloy Designations Temper Designations Superpurity and Commercial Purity Aluminum (lxxx) Al-Si Alloy System
Materials Science and Technology Copyright © WILEY-VCH Verlag GmbH & Co KGaA. Allrightsreserved.
215 217 217 217 217 217 218 219 222 222 223 224 225 227 229 229 229 230 230 230 231 231 231 231 234 235 235 236 237 239 240 240 240 240 241 242
214
5.4.4 5.4.5 5.4.6 5.4.6.1 5.4.6.2 5.4.7 5.4.8 5.4.8.1 5.4.8.2 5.4.8.3 5.4.8.4 5.4.9 5.4.9.1 5.4.9.2 5.4.10 5.4.10.1 5.4.10.2 5.4.10.3 5.4.11 5.4.11.1 5.4.11.2 5.4.11.3 5.4.11.4 5.4.11.5 5.4.11.6 5.4.12 5.5 5.5.1 5.5.1.1 5.5.1.2 5.5.1.3 5.5.2 5.5.3 5.5.3.1 5.5.3.2 5.5.4 5.5.5 5.5.6 5.6 5.6.1 5.6.2 5.7 5.8
5 Aluminum-Based Alloys
Al-Mn-(Mg) Alloy System (3xxx) 242 Al-Mg Alloy System (5xxx) 244 Al-Cu and A l - C u - M g Alloy Systems (2xxx) 246 Al-Cu Alloy System 246 A l - C u - M g Alloy System 247 Al-Mg-Si Alloy System (6xxx) 248 Al-Zn-Mg-(Cu) Alloy System (7250) 250 The Precipitation Sequence 250 Two-Step Aging . . 251 Thermomechanical Treatment (TMT) 252 Commercial Al-Zn-Mg-(Cu) Alloys 254 Al-Li Alloy System 254 The Precipitation Sequence 256 Precipitate Structure and Fracture Toughness 256 Aluminum Powder Metallurgy Products 257 Rapid Solidification 257 Powder Metallurgy (P/M) Alloys Developed by Rapid Solidification (RS). 258 Mechanically Alloyed Aluminum . 259 Aluminum Matrix Composites 259 Maturity and Intrinsic Characteristics of Aluminum Metal Matrix Composites (MMCs) 259 The Processing and Characteristics of Various Metal Matrix Composites (MMCs) 260 Continuous Fiber Reinforced Al Composites 263 Whisker or Short Fiber Reinforced Al Composites 263 Particulate-Reinforced Al Composites 265 Recent Applications of Al Metal Matrix Composites (MMCs) 266 Metallic Glasses and Nanocomposites Based on Aluminum 266 Major Application Fields of Aluminum 267 Packaging and Beverage Cans 267 Flexible Packaging 268 Beverage Cans 268 Closures for Bottles and Collapsible Tubes 269 Building Applications 269 Surface Transportation Industries 269 Automotive Applications 269 Railroad Cars 270 Aerospace Industries 270 Electrical and Electronic Materials 271 Other Applications 272 Concluding Remarks 273 Future Applications and Possible Barriers 273 Aluminum Recycling 274 Acknowledgements 274 References 274
List of Symbols and Abbreviations
215
List of Symbols and Abbreviations b E £ c , f, m / G K Klc KQ m p PD Pz r t T 7^ V{
Burgers vector Young's modulus Young's modulus of composite, fiber, or matrix volume fraction of precipitate shear modulus stress intensity factor critical stress intensity factor (mode I) critical stress intensity factor constant constant driving pressure Zener pressure radius time temperature absolute melting temperature volume fraction of fiber
y 8 e r\c lp Q £(0) a (7C f om, ay T0 Ty
grain-boundary energy elongation or strain strain rate critical particle size interparticle spacing resistivity residual resistivity stress ultimate tensile strength of composite or fiber maximum stress at ultimate tensile strain of fiber yield stress yield strength of the alloy with n o particles, including a solution-hardening term yield strength
CIP CMC CP CTE CVD DC EC ETM FA FCC FILD
cold isostatic pressing ceramic matrix composite commercial purity coefficient of thermal expansion chemical vapor deposition direct chill electrical conductor grade early transition metal Frankford Arsenal face-centered cubic fumeless in-line process
216
FTMT GP HE HIP HVEM IACS I/M ISML ITMT L Ln LNG LT LTM MA MAP MIG MMC PAN PCA PEEK PFZ P/M PMC PSB RA ROM RRA RS RSP SAP SCC SNIF SP SSSS TEM TIG TM TMT UBC UTS VPN WQ YS
5 Aluminum-Based Alloys
final thermomechanical treatment Guinier-Preston hydrogen embrittlement hot isostatic pressing high-voltage electron microscope International Annealed Copper Standard ingot metallurgy Istituto Sperimentale dei Metalli Leggeri intermediate thermomechanical treatment longitudinal lanthanide metal liquid natural gas long transverse late transition metal mechanically alloyed mechanical alloying process metal-inert gas metal matrix composite polyacrylonitrile process control agent poly(ether ether ketone) precipitate-free zone powder metallurgy polymer matrix composite persistent slip band reduction of area rule of mixtures retrogression and reaging rapid solidification; rapidly solidified rapid solidification process sintered aluminum powder stress-corrosion cracking spinning nozzle inert flotation superpurity supersaturated solid solution transmission electron microscopy tungsten-inert gas transition metal thermomechanical treatment used beverage can ultimate tensile strength Vickers pyramid hardness number water-quenching yield stress
5.2 Intrinsic Characteristics of Aluminum
217
5.1 Introduction
5.2.2 Tensile Strength
Metals have played an important role in the development of human civilization. In this development there has not been a metal, apart from steel, as versatile as aluminum, because of its unique intrinsic characteristics (Murakami, 1985, 1991a). Owing to its chemical reactivity, aluminum was not extracted as an industrial metal until 1886, when independent discoveries by Hall in the U.S.A. and Heroult in France led to the establishment of an economic method for its electrolytic extraction. In the relatively short time since then, aluminum production has achieved remarkable growth. In 1921 the annual world production of aluminum reached 200 kilotons (203 kilotonnes), increasing to 1.5 million tons (1.52 million tonnes) by 1950, and 17.5 million tons (17.8 million tonnes) by 1988. Today aluminum is the second most widely used metal, next to steel.
The tensile strength is one of the most important characteristics for any metal. It can be changed over a wide range from very low to fairly high. Although commercially pure aluminum and, particularly, high-purity aluminum are very soft in the fully annealed condition (the latter has a tensile strength of about 50 N/mm 2 ), their strength can be increased to a degree suitable for the intended purpose by alloying and by exploiting versatile strengthening methods such as strain hardening, solidsolution hardening, and precipitation hardening.
5.2 Intrinsic Characteristics of Aluminum Aluminum has unique intrinsic characteristics, which are of interest to scientists as well as engineers. A summary of several main characteristics of aluminum follows, as a basis for further discussions. 5.2.1 Lightness The density of aluminum is only 2.7 g/ cm3, about one third that of steel. Thus aluminum is called a light metal together with a few others, notably magnesium (1.7 g/cm3), beryllium (1.85 g/cm3), and titanium (4.5 g/cm3).
5.2.2.1 Solid-Solution Hardening Solid-solution hardening involves an increase in yield stress produced by elements in solid solution (see Vol. 6, Chap. 5 of this Series). This is most pronounced when the individual solute elements produce elastic distortions in the host lattice which are markedly asymmetric. However, substitutional atoms in a cubic aluminum crystal produce spherically symmetric distortions and hence weak solution hardening. Therefore the practical advantage of solution formation in aluminum is that it may be regarded as a requisite to strain hardening and, in particular, to precipitation hardening. Although most elements can alloy with aluminum, comparatively few elements have sufficient solid solubility to serve as major alloying additions. Of the commonly used elements, only Cu, Zn, Mg, and Si have significant solubilities and play an important role in precipitation hardening. Several transition elements such as Fe, Cr, Mn, and Zr with solubilities below 1 wt.% are used primarily to form intermetallic compounds which give important
218
5 Aluminum-Based Alloys
Table 5-1. Maximum equilibrium and extended (by rapid solidification) solubility of elements in aluminum. Element
Temperature
Equilibrium solubility (mass %)
Extended solubility (mass %)
Ag Cu Zn Mg Li Fe Si Cr Mn Ni Co Mo V Ti Zr Sn Ga Ge
566
55.6 5.67 82.8 14.9 4.0 0.052 1.65 0.77 1.82 0.05 <0.02 0.25 0.6 1.0 0.28 <0.01 20.0 7.2
66.6 34.0 60.0 37.5
548 382 450 600 655 577 660 650 640 660 660 665 665 660 230 30 424
9.7 16.5 11.0 16.8 10.3 10.3 3.5 3.7 3.5 4.9 1.1 82.8 16.8
in industrial practice (see Vol. 6 Chap. 2 of this Series). The work-hardening curves for 1100, 3003 (1.2% Mn), 5050 (1.4% Mg), and 5052 (2.5 % Mg) alloys illustrate that the ultimate tensile strength (UTS) and yield stress (YS) increase rapidly at first and then more gradually, as shown in Fig. 5-1. (The numerical coding system is explained in Sect. 5.4.1.) These increases are obtained at the expense of ductility (as measured by elongation) and also reduce the formability, such as is measured by bending and stretching. However, it should be noted that certain alloys such as 3004 exhibit better drawing properties in the cold worked rather than the annealed condition, and this alloy is used to make thin-walled beverage cans. Work hardening is used extensively to produce strain-hardened tempers from
CO
improvements to the grain structure by controlling recovery and recrystallization. Recently, rapid solidification of aluminum alloys containing the above-mentioned transition elements has been employed to extend the maximum solid solubilities of alloying elements in the primary a-aluminum phase (Table 5-1), and to lead to the reduction or elimination of the second phase, resulting in the microstructural modifications described in Sect. 5.3.5.1.
ROfiP (P R Mg) A
300
CL
L-J
250
5
200
CD
150
I
100
50 0
I
| A
5050 (1.4Mg)-j
3003 (i.2Mn) J
r
I 1100 (Al)
y
CO
a.
I
5.2.2.2 Work Hardening
The properties of a metal are changed by cold working. In particular the mechanical properties such as tensile strength, yield strength, and hardness are generally increased. This increase in the mechanical properties is caused by work or strain hardening, which is not only of interest theoretically, but also of great importance
40
0
^v !
10
20
30 40 50 60 70 80 Reduction by cold rolling. %
90
Figure 5-1. Work-hardening curves for alloys 1100, 3333, 5050, and 5052 (Courtesy of Aluminum Company of America).
5.2 Intrinsic Characteristics of Aluminum
nonheat-treatable alloys. Moreover, working increases the strength achieved through solid-solution and dispersion hardening. In age-hardenable alloys, work hardening complements the strength achieved by precipitation, and may also accelerate the precipitation process. 5,2.2.3 Precipitation Hardening Many important properties of aluminum alloys such as their mechanical strength, toughness, creep, and stress-corrosion cracking are largely affected by the presence of precipitated particles of a second phase. The thermodynamic and some structural aspects of precipitation in binary alloys are set out in Vol. 5, Chap. 4 of this series. The basic requirement for an alloy, necessary for age- (or precipitation) hardening, is a decrease in the solid solubility of one or more of the alloying elements with decreasing temperature. The heat-treatment is normally carried out as follows: (1) solution treatment within the single-phase region, (2) a rapid quench, usually to room temperature, to obtain a supersaturated solid solution (SSSS), and (3) controlled decomposition of the SSSS GP (1) O Al atom
to form a fine, suitable precipitate, usually by aging for convenient times at one or two intermediate temperatures. The decomposition of an SSSS is a complex process which involves several stages and, typically, Guinier-Preston (GP) zones and an intermediate precipitate in addition to the equilibrium phase. The GP zones are coherent, ordered, solute-rich clusters of atoms which often form on one or more atomic planes, e.g., in an Al-4%Cu alloys, as shown in Fig. 5-2. The GP zones are very finely dispersed in the matrix with a density of about 10 17 -10 18 /cm 3 in an Al-4%Cu alloy. The rate of formation is largely influenced by the presence of excess vacancies retained in the matrix by quenching. In some alloys (e.g., Al-Ag), the GP zones are spherical, as illustrated in Vol. 5, Chap. 4, Figs. 4-9 and 4-13 of this Series. The intermediate precipitate is normally much larger in size than a GP zone and is usually partially coherent with the matrix. It has a definite composition and crystal structure which differ only slightly from those of the equilibrium phase. In some alloy systems, such as Al-Zn-Mg, the intermediate phase, 7/'-MgZn2, may be nu-
GP (2) Zone (6") 2 - Layer
Growth
Multilayer
• Cu atom
O
b
r$>
o
P b b o Op i a
o bP o,o o o ooO,bio
C)oo
O(D p 7.68A
4.04A
219
(DP
7.88A
7-8A
Figure 5-2. Structure models proposed for GP(1) and GP(2) (9"-CuAl 2 ) (Yoshida et al., 1982).
220
5 Aluminum-Based Alloys
cleated on the stable GP zones, while in other alloys such phases precipitate heterogeneously at lattice defects such as dislocations, sub-boundaries, etc. The final equilibrium precipitate forms at relatively high temperatures and completely loses coherency with the matrix. This results in little hardening because of a coarse dispersion. Following the discovery of zones by Guinier (1938) and by Preston (1938), the first really comprehensive examination of the structural changes during age-hardening in Al-Cu alloys was carried out by Silcock et al. (1953-54). Figure 5-3 shows that hardening occurs in two stages; the first hardening corresponds to GP(1) formation, and the other to the formation of GP(2) and a fine 0'-CuAl2 intermediate precipitate; further coarsening of the precipitates results in a decrease in the hardness. Gp(l), GP(2), and 6'-CuAl2 denote successive stages in the formation of an intermediate (metastable) phase. In the first stage, there is merely local enhancement of the solute; later, the local crystal structure also begins to change. The crystallography is well explained by Martin (1968); his book also contains reprints of some classic early papers. The strengthening mechanisms in agehardening alloys are explained in terms of obstacles to the motion of dislocations, either by dislocation cutting of fine precipitates such as GP(1) zones, or by dislocations by-passing (or cross-slipping around) widely-spaced precipitates, as illustrated in Fig. 5-4. With respect to the former, the resistance to dislocation motion arises from a number of causes as follows: the dislocation may encounter adverse stress fields due to the atomic volume of the precipitate being different from that of the matrix. Passage of the dislocation through the precipitate may disorder the slip plane
at the interface or within the precipitate. The difference in the stacking fault energy of the matrix and of the precipitate must be taken into account. The yield strength is given by the following equation.
x=cfmrv
(5-1)
where i y is the yield strength, / is the volume fraction of precipitate, r is the radius of the precipitate, and m and p are constants. Once the GP zones are cut, dislocations continue to pass through the GP zones on the same slip planes, and further work hardening thereafter is comparatively small. Thus deformation tends to become localized on a few active slip planes, so that intensive slip bands develop. This microstructure may be deleterious with respect to other properties such as fatigue and stress-corrosion cracking. On the other hand, when precipitate particles are too strong to be cut, the dislocations may be forced to loop around or cross-slip over them by a mechanism first proposed by Orowan (1948), as shown in Fig. 5.4 (see also Vol. 6, Chap. 7, Sect. 7.2.1.2 of this Series). The yield strength i y (in this case) is given by Ty = To
+ 0.8GZ>/Ap
(5-2)
where T 0 is the yield strength of the alloy with no particle including a solution-hardening term, G is the shear modulus, b is the Burgers vector, and Ap is the interparticle spacing. The yield strength is low at the beginning, but the rate of work hardening becomes higher with increasing deformation, because back pressure on the dislocation sources increases on increasing the number of dislocation loops left around the particles. Thus the plastic deformation tends to be spread more uniformly throughout the grains. The hardening reaches a maximum and then decreases with increasing interparticle spacing as a result of
5.2 intrinsic Characteristics of Aluminum
120
Structure GP [1] • 6 © 4 %Cu GP [2] ©3%Cu o4-5%Cu
^S' If
•inn ~Z_ ''""•' Q_
>
y
80
x
£ 500 c o SS^4 % Cu Ann
Jr jr
?/Zone diameter
^^
40 0.1 D
1D 10 D Aging time, days
Dislocation cutting of precipitates
Interaction between dislocation and precipitates Kinds of precipitates and interparticle spacing
(V
s
4UU
o<
300
CD 0)
S»3 % Cu
—-^ If .<•"
60
600
IL'*'0'>\ 41/2 %Cu Ifcfii**
2 % Cu
iarr
140
221
200 0)
o - 100
N
Figure 5-3. Structures and zone diameters of aluminum-copper alloys (Silcock etal., 1953-54).
100 D
Dislocation by-passing (or cross-slipping) by the Orowan mechanism
/ GP zone, coherent
A p <~10" 5 cm small
noncoherent or semicoherent precipitates
Ap >~10"5cm large
Work-hardening rate 6e
elong.
ay(7")oc
A-T2
large
elong.
<x y (7)oc G(T) small
Temperature dependence of yield stress temp. T •
temp. T -
Figure 5-4. Strengthening mechanisms and their characteristics for age-hardened aluminum alloys.
222
5 Aluminum-Based Alloys
coarsening of the particles. This is the situation with over-aged alloys, and thus the typical age-hardening curves proceed as shown in Fig. 5-3. The optimum situation may be realized if the precipitates can resist cutting by dislocations and are too close to permit by-passing by dislocations. In such a case, a high level of precipitate strengthening as well as work hardening is expected, because dislocation motion is only possibly by a process such a cross-slip. Thus a duplex aging treatment carried out below and above the GP solvus has been developed to obtain a very fine dispersion of the intermediate precipitates in some commercial alloys, such as 7000 series alloys. The other possibility may be realized by obtaining a bimodal precipitate structure, which consists of fine, closely-spaced precipitates to increase the yield strength as well as larger precipitates to raise the rate of work hardening and simultaneously to cause more uniform plastic deformation. 5.2.3 Precipitate Structures and Technological Properties
In commercial age-hardening aluminum alloys, there are generally three kinds of precipitate particles which control the technological properties: (i) Coarse intermetallic compounds, often containing Fe and Si of 0.5-1.0 jim in diameter, which form during solidification or during subsequent homogenization and hot working. They are either virtually insoluble compounds, e.g., Al 6 (Mn,Fe), Al 3 Fe, a-Al (Fe,Mn) Si, and Al 7 Cu 2 Fe, or relatively soluble compounds, e.g., CuAl 2 , Mg2Si, and S (Al2CuMg). These particles tend to be aligned as stringers in as-fabricated products. (ii) Smaller particles or dispersoids of 0.05-0.5 jim in diameter, which are intermetallic compounds containing the transi-
tion elements Cr, Mn, or Zr. They include Al 20 Cu 2 Mn 3 , Al 12 Mn 2 Cr, or Al3Zr particles, which primarily inhibit recovery, recrystallization, and grain growth (see Vol. 15, Chap. 9, Sect. 9.8 of this Series). (iii) Fine precipitates, up to 0.01 jim in diameter, which precipitate during aging heat-treatments and promote strengthening, as described in Sect. 5.2.2.3. These particles produce a great variety of precipitate structures which exert a large influence on the technological properties of commercial alloys. These important subjects are explained in the following sections (see also Murakami, 1990). 5.2.3.1 Bimodal Precipitate Structures and Grain Refinement
In commercial alloys containing about 1 wt.% of transition elements, a bimodal precipitate structure consisting of a coarse distribution of large particles and a fine dispersion of smaller particles results during heat-treatment, as shown in Fig. 5-5 (Nes, 1980, 1986). The effect of such a bimodal precipitate structure on the recrystallization behavior of deformed alloys has
Coarse particles
Deformation zone
Deformed 'sub-structure
Fine distribute dispersoic
Figure 5-5. Sketch outlining the principal features of the cold-deformed substructure of an alloy with a bimodal precipitate structure (Nes, 1980).
5.2 Intrinsic Characteristics of Aluminum
been studied extensively in recent years (see Vol. 15, Chap. 9, Sect. 9.8.5 of this Series). Small, finely dispersed precipitate particles inhibit recrystallization, and large, coarsely distributed particles promote recrystallization due to deformation zones surrounding the coarse particles. Particle stimulated nucleation has been found to occur above a critical particle size, rjc, which is given by 2y
(5-3)
where y is the grain-boundary energy, P D is the driving pressure for recrystallization due to the stored deformation energy, and Pz is the 'Zener pressure' which is due to the drag from small dispersoid particles. Based on TEM observation of particle stimulated recrystallization, Humphreys (1979, 1980) divided the recrystallization process into three stages: the formation of a nucleus from preexisting subgrains (Fig. 5-6), recrystallization of the deformation zone, and growth of recrystallization beyond the deformation zone. A very fine, recrystallized grain structure can be obtained if the above criteria are satisfied. Warlimont (1977) and Theler and Furrer (1974) studied the recrystallization of commerical 0.7-1.0 wt.%Mn aluminum
Figure 5-6. Recrystallization nucleation at a cluster of SiO2 particles during a nickel in situ HVEM anneal (Humphreys, 1979).
223
alloys at temperatures, where recrystallization and precipitation can occur simultaneously and competitively. For an Al0.7 %Mn alloy at temperatures above 763 K, recrystallization is completed before precipitation begins, and thus a fine grain structure results. However, at temperatures between 623 K and 593 K, recrystallization is greatly retarded, because precipitation becomes dominant. This results in a highly inhomogeneous, coarse grain structure. 5.2.3.2 Precipitate Structures and Fracture Toughness
Control of the precipitate structure has recently been recognized as being very important for improving the fracture toughness, since it could have a great influence on both the initiation and propagation of cracks. Figure 5-7 schematically shows the mechanism for the nucletion of cracks under tensile stress (Rosenstein, 1982). The important metallurgical factors are: (a) In the underaged condition, GP zones are easily cut through by dislocations, and thus deformation results in localization of the strain in bands across the grains. These deformation bands lead to dislocation pile-ups at the grain boundary, and wedgeopening grain-boundary failure results, (b) In the overaged condition, a precipitate-free zone (PFZ) results on both sides of the grain boundary. The PFZ is soft and thus deformation results in the localization of strain in the PFZ and grain-boundary region, which serves as a nucleation site for microvoid formation; this is followed by microvoid coalescence. Thus low-ductility grain-boundary failure results, (c) Crack nucleation and propagation proceed by the cracking of large particles or by decohesion at the interface between second particles and the matrix.
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5 Aluminum-Based Alloys
Figure 5-7. Schematic drawing showing the crack nucleation mechanism: (a) The localization of strain into bands across the grain; (b) crack occurrence at the triple point of the grain boundary, and grain-boundary failure by preferential deformation of the weak PFZ (precipitate-free zone) region; and (c) crack nucleation and propagation proceeding by the cracking of large particles or by decohesion at the interface, T is stress.
In the 1950s metallurgists at ALCOA developed a high-strength Al-Cu-Li alloy (2020). However, this alloy had low ductility and fracture toughness after the maximum strength temper. ("Temper", in general, denotes a condition generated by a mixture of heat-treatment and plastic deformation. For details, see Sect. 5.4.1). These limitations led to its withdrawal from commercial consideration in 1969. Recently, with the progress in fundamental research on fracture mechanics, improvements in the ductility of Al-Li alloys could be obtained by controlling the microstructures of the alloys. Additions of small amounts of Zr to Al-Li-Cu and A l Li-Cu-Mg alloys, and suitable thermomechanical treatments result in beneficial microstructures for increasing the fracture toughness. Zr precipitates coherently and homogeneously in the form of oc'-Al3Zr particles, which control grain size and shape. Moreover, a'-Al3Zr precipitates are known to act as nucleation sites for 8'Al3Li, and thus 5'-Al3Li particles form complete envelopes surrounding the oc'Al 3 Zr particles, and lead to the formation of composite precipitates which cannot be sheared by moving dislocations (Gayle and Vander Sande, 1984).
A good, recent overview on the present state of development of Al-Li-based alloys is by Grimes (1990). 5.2.3.3 Precipitate Structures and Fatigue Strength
It is well-known that alternating stresses impair the strength of materials considerably. While steel has a fatigue strength of nearly half its tensile strength, the fatigue strength of high-strength, age-hardened aluminum alloys is only about one-third or less of the alloy's tensile strength. This feature is illustrated in Fig. 5-8, which shows the relation between the endurance limit (5 x 108 cycles) and the tensile strength for various alloys (Polmear, 1989). In order to improve the fatigue strength of age-hardened alloys, it is necessary to improve their resistance to crack initiation and propagation. Crack initiation is known to normally occur at the surface, where strain is localized due to the presence of notches, corrosion pits, persistent slip bands (PSBs), in wich extrusions and intrusions may form, or PFZs at the grain boundaries (see Vol.6, Chap. 11, Sect. 11.2.2 of this Series). The important effects of precipitate microstructures on the fa-
5.2 Intrinsic Characteristics of Aluminum 250
I
-200 0
225
x aged aluminum alloys O nonheat-treatable aluminum alloys •
magnesium alloys
A steels
o 150 x 10
100
I 50
Figure 5-8. Fatigue ratio (endurance limit/tensile strength) for aluminum alloys and other alloys (Polmear, 1989).
HI
100
200
300 400 500 Tensile strength (MPa)
tigue behavior have been described in detail by Starke and Lutjering (1979). The poor fatigue properties of age-hardened aluminum alloys may be attributed to two factors, i.e., the metastable precipitates being shearable by the dislocations moving within the grains and the presence of PFZs at the grain boundaries. When shearable precipitates are present in the alloys, narrow PSBs are formed and crack nucleation results, as mentioned in the previous section. Where PFZs are present at grain boundaries, localized plastic deformation may occur in these soft regions with resultant crack nucleation at the grain boundaries. In order to improve the fatigue properties, the influence of both types of precipitate structure must be avoided. As for PFZs, it may be effective to introduce obstacles which resist localized shear by the following two means. The first is to introduce nonshearable particles by powder metallurgy (P/M), or moderately large, intermediate precipitates by slight overaging. The fine grain size may also be effective in preventing the formation of PSBs, since it reduces the slip length and thus hinders the formation of large pile-
600
700
ups of dislocations at the grain boundaries. Secondarily, in order to avoid the drawbacks of PFZs, reducing the width of the PFZs by two-stage aging may be effective, or changing the shape of the grains by thermomechanical treatment, which produces large flat grain boundaries. A decrease in the grain size may also be effective, as in the previous case. 5.2.3.4 Precipitate Structures and Fine-Grain Superplasticity
On the basis of existing knowledge on superplastic alloys, it is evident that a finegrained material must be produced (with grains of the order of a micrometer in diameter) which is not susceptible to rapid grain coarsening at temperatures where superplastic flow, i.e., mainly grainboundary sliding, occurs. The first common superplastic aluminum alloys were of eutectic or eutectoid composition, because both recrystallization and grain growth tend to be restricted by the presence of small, separate grains of two phases. However, it is considered that if the grain size of an alloy remains very fine during high-
226
5 Aluminum-Based Alloys
temperature deformation, owing to the presence of a large amount of dispersed second phase remaining undissolved, which inhibits recrystallization, a nominally single-phase alloy could behave superplastically when deformed at suitable strain rates and temperatures. Since Cr, Mn, and particularly Zr can have a marked effect by inhibiting the recrystallization of aluminum alloys, fine-grained, nominally single-phase alloys can be prepared by the control of precipitate structures by static as well as by dynamic recrystallization. Figure 5-9 a shows a representative optical microstructures of the Al-5.8Mg0.37 Zr-0.07 Cr-0.16 Mn alloy, which was
Figure 5-9. (a) Optical microstructure of Al-5-8 Mg -0.37 Zr-0.07 Cr-0.16 Mn alloy obtained after 200% elongation under tensile testing at 793 K using g = 8.3xlO" 4 s~ 1 , and (b) TEM micrograph showing a very fine distributioin of oc'-Al3Zr particles (Matsukietal., 1976).
elongated by 200 % at an initial strain rate of 8.3xlO~ 4 s" 1 , and showed no grain growth. Figure 5-9 b shows a TEM micrograph depicting the internal structure of the specimens, which shows the presence of fine particles of ~ 30 nm in diameter with contrast due to coherency strain as well as some dislocations which are pinned by these fine particles (Matsuki et al., 1976). Grimes et al. (1976) have developed the Al-6Cu-0.4Zr (Supral 100) and A l 6Cu-0.4Zr-0.2Mg-0.1Ge (Supral 220) superplastic alloys. These alloys have a partially polygonized, unrecrystallized structure before superplastic forming, which transforms into a very fine, recrystallized grain structure during 50-100% deformation at high temperatures. These very fine grain structures are believed to be produced by dynamic recrystallization of the subgrain structure and stabilized by very fine oc'-Al3Zr particles. Higashi et al. (1986) supported this idea and proposed that a more fine-grained structure could be obtained more easily using dynamic, continuous recrystallization at high temperatures, and they succeeded in increasing the strain rate, so reducing void formation. Wert etal. (1981) reported on a thermomechanical process for grain refinement in the precipitation-hardening 7075 aluminum alloy, which has substantial potential for superplastic forming operations for mass production. The process includes the creation of preferential nucleation sites where grains recrystallize around the localized deformation zones associated with T|'MgZn 2 particles which are larger than the critical size, as well as the Zener drag by precipitated, fine oc'-Al3Zr particles. A four-step thermomechanical process sequence has been devised for grain refinement of the 7075 alloy, and is as follows:
5.2 Intrinsic Characteristics of Aluminum
(i) Solution treatment: 755 K for 10.8 ks (i.e., 482 °C for 3 h), then water-quenching (WQ); (ii) overaging: 673 K for 28.8 ks (i.e., 400°C for 8 h), then WQ; (iii) warm rolling at 473 K; 90% reduction using 5 10 passes; and (iv) recrystallization; 755 K for 1.8 ks (i.e., 482 °C for 30min), then WQ. The overaging step produces particles larger than approximately 0.75 ^m in diameter at a high enough density for there to be about ten in each approximately 10 jim recrystallized grain. 5.2.3.5 Precipitate Structures and Stress-Corrosion Cracking
Some aluminum alloys are known to rupture mainly along grain boundaries under the combined action of static or dynamic tensile stress and a specific corrosion environment. This phenomenon is called stress-corrosion cracking (SCC). (SCC, as it pertains to steels, is discussed in Vol. 7, Chap. 12, Sect. 12.15 of this Series). SCC is limited to Al-Cu-Mg alloys, A l Mg alloys containing more than 3 % Mg, and Al-Zn-Mg-Cu alloys, and rarely occurs in Al-Mg-Si alloys, while Al-Mn and Al-Mg alloys that contain less than 3% Mg show no sensitivity to SCC. However, at present the most widely used highstrength aluminum alloys are generally resistant to SCC under ordinary service conditions when treated using the proper temper conditions, although the usefulness of the strongest 7075-type aluminum alloys is often limited by their susceptibility to SCC, because severe failures take place in peak-aged condition. This susceptibility can be eliminated by overaging (i.e., T7type temper), but at a loss of about 15% in strength. The mechanisms of SCC in Al-Zn-Mg alloys have been studied extensively from the viewpoint of precipitation phenomena
227
(Murakami, 1981). The early view of the SCC mechanism in Al-Zn-Mg alloys, was that cracking occurred mainly by anodic dissolution. There is a precipitate-free zone (PFZ) on both sides of the grain boundary in peak-aged Al-Zn-Mg alloys. Cracking occurs at PFZs by mechanochemical dissolution under the conditions of stress concentration and in corrosive environments. Therefore the main factors influencing SCC have been considered to be the width of the PFZ, and the size and distribution of precipitates within the grains as well as on the grain boundaries. However, some studies (Gruhl, 1963, Scamans et al., 1976) have indicated that crack propagation may occur as a result of hydrogen embrittlement (HE), as opposed to anodic dissolution. Observations of reversible pre-exposure embrittlement, of cracking in humid air containing too little moisture for effective dissolution, of greater susceptibility in mode I than in mode III loading, of cathodic charging and of cracking with no discernible grain-boundary precipitate dissolution on the SCC fracture surface, have together definitely proved that HE is involved in the SCC mechanism (Gest and Troiano, 1974; Scamans, 1980). Since the SCC of Al-Zn-Mg alloys is mostly intergranlar, numerous studies have focused on grain-boundary segregation. It has been shown that the Mg and Zn concentration profiles across the PFZs and the grain boundaries vary greatly, depending on the heat-treatment and the quenching rate; this was done using electron energy loss measurements and X-ray microanalysis in a transmission electron microscope (Doig and Edington, 1975). From experimental results exploiting the extreme surface sensitivity of Auger electron spectroscopy to obtain in situ segregation information on the actual grain boundary in the specimens of various Al-Zn-Mg and Al-
228
5 Aluminum-Based Alloys
Zn-Mg-Cu alloys, Chen et al. (1980) concluded that both Mg and Zn segregate at the boundaries, with the Zn largely bound in MgZn 2 precipitates, but with a significant portion of Mg atoms existing "freely" on the boundaries. The presence of this free Mg in the region where cracking occurs and the great affinity of Mg for hydrogen are together believed to give strong support to an Mg-H interaction mechanism for SCC (Scamans, 1980, Viswanadham et al., 1980). Rapidly solidified aluminum powder metallurgy (P/M) alloys, 7091 and 7090 (which have 0.4 or 0.6 mass% Co added to 7075, respectively), are known to show improved combinations of strength and SCC resistance. Rapid solidification produces two sizes of Co 2 Al 9 particles, smaller particles within grains and larger particles on the grain boundaries. The SCC susceptibility of P/M 7091 and 7090 rapidly solidified (RS) alloys was investigated by Pickens and Christodoulou (1987). Their explanation of the SCC mechanism in these alloys is summarized as follows: The surface oxide film is most readily penetrated by aqueous environments in the regions near large particles. Once the oxide film is penetrated and a microcrack forms, dissolution will occur in the grain-boundary region in the presence of water. The cathodic half-reaction of the dissolution, at the pH levels shown to exist at the crak tip (~ 3.5), is hydrogen ion reduction, i.e., H + + e ~ -* H, which can lead to adsorption of hydrogen on the grain boundaries. Hydrogen atoms then diffuse along the grain boundaries ahead of the crack tip. This diffusion may be enhanced by the driving force of the triaxial stress field at the crack tip. In addition, grain-boundary diffusion of hydrogen is increased by the free Mg segregation observed there. Furthermore, the free Mg may cause the hydrogen to
remain on the boundary. In any event, when a critical hydrogen concentration is reached, fracture takes place on the boundaries at stresses lower than those that cause fracture without any environmental influence. However, exposed Co 2 Al 9 particles on the grain boundaries of the RS-P/M alloys serve as sites for hydrogen recombination and lose hydrogen gas to the atmosphere, thereby reducing the hydrogen concentration absorbed at the grain boundaries, and so improving the resistance to SCC. An excellent heat-treatment known as retrogression and reaging (RRA) has been proposed by Kaneko (1980), Rajan etal. (1982), Islam and Wallace (1983), and Danh et al. (1983) for 7000 series alloys to obtain properties equivalent to those obtained by the T73 temper together with the T6 strength level. RRA is a two-step treatment, as shown in Fig. 5-10. During the first step, the 7075-T6 alloy, for example, is retrogressed at a temperature between 478 K (205 °C) and 533 K (260 °C) for a short time, water-quenched, and then, during the second step, reaged at the original temperature of 393 K (120 °C). In the initial step, retrogression causes the partial dissolution of the pre-existing GP zones, but not of coarse precipitates at the grain boundaries. The remaining GP zones then act as nucleation sites for rjf particles. The dissolution of GP zones, furthermore, enriches the matrix in Zn and Mg, which in turn promotes nucleation and growth of the rj' phase. In the second step, both the GP zones and the i\' precipitates can nucleate and grow. The coherent fine precipitates help to raise the strength level, while the coarse grain-boundary precipitates improve the SCC resistance.
5.2 Intrinsic Characteristics of Aluminum o x • •
Retrogressed at 493 K and re-aged Retrogressed at 493 K Retrogressed at 473 K and re-aged Retrogressed at 473 K
229
Retrogressed and re-aged T6 yield
0.6 1.2 1.8 Retrogression time, t/ks
3.6
7.2 10.8
5.2.4 Physical Properties Relating to Commercial Applications 5.2.4.1 Electrical Resistivity
The measured value of the electrical resistivity of "four-nines" (99.99%) purity aluminum at 20 °C has been reported to be 2.6548 jiQ cm, which corresponds to a conductivity of 64.94% of the International Annealed Copper Standard (IACS). Therefore an aluminum conductor may have more than 60 % of the conductivity of a copper conductor, and thus when two conductors of equal length and weight are compared, the aluminum conductor can transport twice as much current as the copper conductor. The electrical resistivity generally increases with increasing solute in pure aluminum, and at very low temperatures (below 100 K) becomes highly sensitive to the degree of purity. Therefore the ratio of the resistivity at 293 K to that at 4.2 K is called the "residual resistivity" and is used as a measure of the purity. A measured ratio as high as 30 000 has been obtained for five-nine purity material. The temperature dependence of the resistivity of aluminum at temperatures below 100 K is expressed as (5-4)
Figure 5-10. Changes in yield strength during retrogression at 473 and 493 K, and after retrogression plus reaging in the RRA treatment for the improvement of the SCC resistance of 7000 alloys (Danh et al., 1983).
where Q (0) is the residual resistivity, T is the temperature in Kelvin, and A and B are constants that can be determined for each purity. The T2 and T5 terms arise from electron-electron scattering and electronphoton scattering, respectively. The resistivity of all high-purity metals generally decreases rapidly and steadily with decreasing temperature. In the case of aluminum, this decrease is particularly large and exceeds that of copper. Hence, below 62 K, high-purity aluminum has a lower electrical resistivity than high-purity copper. Furthermore, the electrical conductivity of aluminum at very low temperatures is less harmfully affected by strong magnetic fields than that of copper. 5.2.4.2 Thermal Conductivity
Aluminum is known to be a good heat conductor, and its thermal conductivity is about half that of copper, but three times that of iron and twelve times that of stainless steel. The true thermal conductivity of fully annealed, high purity aluminum of better than four-nines is relatively insensitive to the impurity level at moderate to high temperatures of more than about 100 K, but below this temperature it becomes highly sensitive to the impurity level measured by the residual resistivity. The
230
5 Aluminum-Based Alloys
measured value of the thermal conductivity of fully annealed, high-purity aluminum with a £(0) of 5.95 x 10~ 4 \xQcm is reported to be 2.36 W c m " 1 K" 1 at 273.2 K. 5.2.4.3 Other Physical Properties Aluminum and its alloys are slightly paramagnetic materials. The small amounts of iron that are usually present as an impurity have little effect because iron usually forms a second phase consisting of an aluminum-iron intermetallic compound such as Al 3 Fe, which is paramagnetic, instead of forming ferromagnetic particles, as happens in copper and its alloys. Thus aluminum and its alloys are usually accepted as nonmagnetic materials. Aluminum has a very low cross section for thermal neutrons of 0.23 barn (0.23 x 10~ 28 m 2 ). Hence aluminum can be used in the nuclear field as a structural material which is virtually transparent to neutrons, like Mg [0.059 barn (0.059 x 10~ 28 m 2 )], Be [0.009 barn (0.009 x 10~ 28 m 2 )], and Zr [0.18 barn (0.18 x 10~ 28 m2)]. Aluminum alloys are often used as components of dispersion fuels (see Vol. 10 A, Chap. 2, of this Series). 5.2.5 Corrosion Characteristics 5.2.5.1 Surface Oxide Film Aluminum should have a low corrosion resistance principally due to the fact that it is very active thermodynamically. However, most aluminum alloys actually have good corrosion resistance due to the presence of a thin, adherent film of compact aluminum oxide on the surface. If a fresh aluminum surface is exposed to air or water, a surface oxide film forms and grows very quickly. This protective film is generally stable in aqueous solutions of pH between 4.5 and 8.5, whereas it is soluble in
strong acids or alkalis leading to rapid attack, with the exception of concentrated nitric acid, glacial acetic acid, and concentrated ammonium hydroxide. The surface oxide layer formed in dry oxygen is very thin, e.g., 2.5-3.0 nm at room temperature. At higher temperatures the oxide film grows more rapidly; the kinetics are expressed by a logarithmic time law for low to moderate temperatures. Higher humidities increase the oxide film thickness, and at room temperature and 100% relative humidity about twice as much oxide is formed as in dry oxygen. The initial corrosion product in aqueous solutions is generally aluminum hydride, which changes with time to hydrated aluminum oxides. The main difference between this film and that formed in air is that it is less adherent and hence it is less protective. At higher temperatures a more complex oxide film is formed on the aluminum alloy containing Mg and Cu, with kinetics that are no longer controlled by simple time laws. The anodizing process using various chemical and electrochemical treatments can produce a much thicker surface oxide film, e.g., 10-20 jim thick, which improves the resistance of aluminum and its alloys to corrosion. In the anodizing process, the aluminum material acts as the anode in an electrolyte such as an aqueous solution containing 15% H 2 SO 4 , producing a porous A12O3 film, as shown in Fig. 5-11. This porous film is subsequently sealed by treatment in pressured water vapor. In addition, color can be introduced into the porous film or even developed from the alloy itself during the anodizing process. This process is applied to architectural and ornamental products.
5.3 Advantages in Manufacturing and Fabrication Processes
Figure 5-11. Keller-Hunter-Robinson model for the fine structure of the oxide film anodized at 120 V in a phosphoric acid bath.
5.2.5.2 Contact Corrosion Behavior with Other Metals
Aluminum and its alloys are known to be attacked by contact with some other metals. The electrode potential of aluminum with respect to other metals has to be considered when a galvanic reaction may arise. For example, magnesium with an electrode potential of —1.73 V is more negative than aluminum with —0.85V, whereas mild steel is more positive with a value of-0.58 V. Thus the sacrificial attack of aluminum and its alloys may occur when they are in contact with mild steel. 5.2.5.3 Influence of Alloying Elements on Corrosion Alloying elements and impurities may influence the corrosion resistance to varying degrees, depending upon whether they
231
are present in solid solution or as microconstituents. Magnesium in solid solution confers a relatively high resistance to corrosion by sea-water and alkaline solutions. Silicon, chromium, and zinc in solid solution in aluminum have only a minor effect on the corrosion resistance, whereas copper reduces the corrosion resistance of aluminum more than any other element. However, small amounts of copper, e.g., 0.05-0.2 wt.%, tend to reduce pitting attack due to inducing general corrosion, and hence retard perforation by pitting, although the overall weight loss becomes larger. Microconstituents can occasionally create the most serious problems with electrochemical corrosion, leading to localized attack such as pitting, and intergranular and exfoliation corrosion. The microconstituents in the grain boundaries have a significant relationship with SCC, as described in Sect. 5.2.3.5. The rate of general corrosion of highpurity aluminum is much less than that of the commercial purity grades, which is attributed to the smaller size and number of the cathodic constituents of iron and silicon compounds. Manganese forms MnAl 6 with aluminum, which has almost the same electrode potential as aluminum. Magnesium contents above the solid solubility in aluminum tend to form the strongly anodic phase Mg 5 Al 8 ; this precipitates in grain boundaries and promotes intercrystalline attack.
5.3 Advantages in Manufacturing and Fabrication Processes 5.3.1 Melting and Casting Operations
An obvious advantage of aluminum is its comparatively low melting point (pure Al: 660 °C), as well as its easy melting and
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5 Aluminum-Based Alloys
pouring in air without any protective flux covering or inert gas atmosphere, since the oxide surface skin that forms on the molten metal prevents excessive oxidation. The low melting point of aluminum ensures that the furnaces and ladles have a fairly long life, even when using relatively simple thermal insulator and ladle linings. Melting can easily be achieved using oil, gas, or electricity and, moreover, handling is also relatively easy because of the low density of aluminum. The production of semifinished products such as plate, sheet, strip, extrusions and forgings generally starts by rolling slab or extrusion billets which are manufactured by the direct chill continuous casting (DC) process. As shown in Fig. 5-12, molten metal is poured into a watercooled, frame-shaped metal die with a bottom block, which can be lowered at a constant rate of about 5-10cm/min; equiva-
loader
7 cooling water
Figure 5-12. Schematic diagram of DC (direct chill) casting.
lent to the rate of molten metal pouring. Several slabs or billets can now be cast simultaneously in the same unit and the weight of the rolling slabs has gradually been increased to 15-18 tons. On the other hand, both strip and plate with a thickness of about 2-5 mm have been produced by the continuous casting processes shown in Fig. 5-13 (Slevolden, 1974). These continuous strip-casting processes make it possible to produce strip products direct from molten metal without having to invest in an expensive, large hot mill. Hence this process is particularly beneficial in cases where capital for industrial investment is limited as for a small scale enterprise, or where urgent demands for increased strip production do not allow time for the construction of a large hot mill. DC casting produces a fine-grained structure compared to mold casting and, moreover, strip casting produces finer microstructures. In the latter process, the metal cast in the solidification region flows into the mold formed by the water-cooled rotary cylinder, caterpillar tracks, or endless steel belts, as shown in Fig. 5-13. Hence strip casting leads to rapid solidification of the melt and rapid cooling to temperatures well below the solidus line. Alloying elements present in amounts above their solid solubility limits, particularly in the case of transition elements such as iron, manganese, chromium, zirconium, and titanium, are either completely or else at least partially suppressed into supersaturated solid solution. This SSSS is decomposed during hot working or subsequent heat-treatments to form very fine, homogeneous dispersoids which play an important role in recovery and recrystallization, resulting in an improvement in the mechanical properties, as mentioned in the previous section.
233
5.3 Advantages in Manufacturing and Fabrication Processes Hazelett steel m ) belts
N
/water' cooled \ v wheel /
water- ^ \ cooled J Awheel A
solid
Spidem, Southwire, Mann
Properzi
liquid reservoir liquid solid Hunter-Douglas
Hunter Engineering
Horizontal sheet casting
Figure 5-13. Diagrammatic sketches of common continuous casting processes based on the moving mold principle (Slevolden, 1974).
As-melted molten metals generally include gaseous elements (hydrogen), oxides, intermetallics, inclusions, and alkali and alkaline-earth elements such as sodium, lithium, calcium, etc. Since these inclusions impair the hot and cold formability as well as the mechanical and chemical properties, they have to be removed as completely as possible from the molten metals before casting into slabs and billets. Several efficient technologies have been developed in the following domains: (i) Degassing processes: chlorine gas injection, flux treatment; (ii) filtering of oxides, intermetallic, or other inclusions: ceramic tube filtering, alumina ball filtering; and (iii) in-line systems: FILD (fumeless in-line process), SNIF (spinning nozzle inert flotation) (Fig. 5-14; Kimzey, 1978). As for aluminum castings, small aluminum foundries can be set up with little capital expenditure, owing to the ease of melting and casting aluminum. There are many casting techniques, such as sand casting, resin-mold and permanent-mold
casting, pressure-die casting, and other processes, which can produce cast products of sound quality. The quality has been improved substantially by better liquid metal treatment and also by improved techniques. Premium-quality castings can be produced by adopting appropriate, special-casting techniques.
melt exit
heater rotating nozzle rotating nozzle
graphite pipe
Figure 5-14. Schematic diagram of the SNIF process (Kimzey, 1978).
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5 Aluminum-Based Alloys
5.3.2 Versatile Deformation Techniques
Aluminum is intrinsically readily workable and easy to form by versatile deformation techniques, owing to its soft, facecentered cubic (FCC) crystal structure. Industrial working operations are divided into either hot or cold working, respectively, for temperatures above and below 0.5 Tm (where Tm is the absolute melting temperature). Hot working is generally carried out at temperatures between 350 and 530 °C, where dynamic recovery or static recrystallization occur during or after deformation. Cold deformation is performed at room temperature, although the material itself may reach temperatures up to about 150°C due to the stored energy that it accumulates during the process. The final stages of mechanical working are generally carried out at room temperature with the purpose of strengthening and to improve the surface quality, and in order to attain precisely determined dimensions, such as sheet thickness. Hot deformation allows not only a high degree of deformation, but also serves as a preliminary step to subsequent cold working since the as-cast, inhomogeneous structure is destroyed suitably for cold working. The most popular hot-working processes for wrought materials are the rolling of ingots into the semiproducts of sheet and plate, and the extrusion of billets from which the economic and direct production of shaped materials with various complex cross sections can be carried out, as shown in Fig. 5-15. After hot rolling, sheet or coiled sheet is normally cold rolled with an intermediate anneal. The as-cold-rolled or susequently annealed sheets are intended to have appropriate microstructures (grain structure, dislocation substructure, and texture) for the following metal-forming processes.
i\7\r Figure 5-15. Various complex extruded sections of high-strength aluminum alloys.
Texture is synonymous with preferred orientation, in which the individual grains in a polycrystalline material are not randomly orientated, but nearly all have the same orientation. Thus the metal-forming of a sheet with texture can result in defects, such as earing in cup samples. Earing may lead to production problems due to trouble in removing the rim of deep-drawn products. Earing may be minimized by careful control of the rolling and annealing schedules. For sheet materials, rerolling, roll-forming, pressing, deep drawing, and spinning, etc. are commonly used to manufacture the finished articles. (For further details, see Vol. 15, Chap. 10, Sects. 10.3 and 10.4 of this Series.) Extruded materials can afford some very interesting and promising applications of aluminum structural members in the building and transportation industries. In this field extrusions are further deformed by drawing, die-forging, stretching, bending, and electro- or magnetoforming at room temperature or at a slightly higher temperature.
5.3 Advantages in Manufacturing and Fabrication Processes
Die-forging and impact extrusion are very suitable and economic manufacturing processes for the mass production of small and medium-sized articles. Die-forging of rod sections in several stages gives the near-net shape of the finished piece, which needs little machining by upper and lower dies. The impact extrusion of slugs, in some cases followed by an ironing operation, has provided a large market for sheet materials in the beverage can industriy. The commercial importance of finegrained, superplastic forming is best recognized for higher strength alloys where conventional forming methods do not satisfy these requirements. 5.3.3 Good Machinability
Aluminum offers rapid and economical machinability, particularly for normal casting alloys and for rods manufactured from special free-cutting alloys. The microconstituents present in aluminum alloys considerably affect the machinability. Nonabrasive constituents have a beneficial effect, whereas insoluble, abrasive constituents exert a detrimental effect on tool life and surface quality. Constituents that are insoluble but soft and nonabrasive are beneficial because they assist in chip breakage, and hence constituents such as bismuth and lead are deliberately added to high-strength, free-cutting alloys for processing in high-speed, automatic bar-cutting machines. Alloys containing more than 10% silicon are the most difficult to machine, because hard particles of free silicon cause rapid tool wear. High-strength, age-hardenable alloys containing fairly high concentrations of alloying elements such as copper, silicon, magnesium, and zinc can be machined relatively easily to a good finish with or without cutting fluid, but a cutting fluid is recommended for
235
most operations. The machinability of aluminum alloys is classifed into five groups: A, B, C, D, and E, in increasing order of chip length and in decreasing order of finish quality. Further details can be found in the Metals Handbook (1967). 5.3.4 Versatile Bonding Technologies
Aluminum has enormous potentialities for building large-scale constructions that can resist high stresses, because of its relatively low weight, high specific strength and modulus, high electrical and thermal conductivity, resistance to corrosion, and good appearance, and moreover it can be bonded by various fastening techniques such as riveting, bolting, and also by various welding techniques, the technologies of which have been greatly improved in recent years together with the development of advanced materials of high specific strength and modulus. The commonly used aluminum alloys for fasteners are as follows: 1100 for cold-formed rivets, 6061T6 and 2024-T4 for machine screw nuts, cold-formed rivets and bolts, and 2011-T3 for screws and bolts. It has been proved for light-metal railway vehicles that aluminum is not only profitable in service but also in fabrication when use is made of jigs that permit efficient production and guarantee a dimensional accuracy hitherto unattainable. As for welding, inert gas shielding arc welding is widely adopted, because this process needs only moderate skill and automation is easily applicable. This process is particularly useful for welding aluminum and includes two kinds of process, as follows. In metal-inert gas (MIG) arc welding a bare wire is continuously fed into the weld from a large spool, as shown in Fig. 5-16. The wire acts as the electrode and filler in the joint. An inert gas, usually
236
5 Aluminum-Based Alloys feeding wire electrode wire
- electrode chip gas nozzle
arc
melt pool
welded metal
Figure 5-16. Schematic diagram of MIG arc welding.
argon or helium, is also fed into the torch. Tungsten-inert gas (TIG) arc welding uses a "nonconsumbale" tungsten electrode. A separate filler rod is melted into the joint when required, but the electrode itself is not used as a filler (Fig. 5-17). Electron-beam welding is a welding process in which the heat is produced by bombardment with a dense stream of highvelocity electrons, which are produced in a high vacuum better than 10 " 4 Torr (0.0133 N/m 2 ) by an electron gun. The major advantages of this process include: (1)
W electrode i
the ability to make welds that are deeper, narrower, and less tapered than the best gas-tungsten arc welds, with a total heat input that is much less than that used in arc welding; (2) superior control over penetration and weld dimensions and properties; and (3) freedom from inclusions such as oxides and nitrides. Aluminum-lithium alloys are not easy to weld, and it is interesting that a specially designed Al-Cu-Li alloy, Weldalite 049, has been designed for ready weldability (Pickens, 1990). Brazing is also a widely utilized joining process for aluminum goods such as automobile condensers. The brazing of aluminum alloys was made possible by the discovery of fluxes that break the oxide film on aluminum without causing any damage, and by the development of filler metals with suitable melting ranges and other properties. The filler metal is distributed between the closely fitted surfaces of the joint by capillary action. A recent development is the use of brazing foils (e.g, Al-Si eutectic composition) made by rapid solidification (Rabiukin and Liebermann, 1993). Other joining methods, e.g., diffusion bonding, electrical resistance welding, flashbutt welding, ultrasonic welding, and others, are used with aluminum and are vital for certain applications. 5.3.5 Advanced Processing Technologies
:*| gas nozzle
welding bar welded metal
melt pool Figure 5-17. Schematic diagram of TIG arc welding.
Until recently, a limited improvement in the microstructure-related properties affecting the performance of conventional aluminum and its alloys continued to be obtained by impurity control, minor element additions, and improved heat-treatment, as well as by thermomechanical processing. In addition to those mentioned above, new secondary manufacturing techniques,
5.3 Advantages in Manufacturing and Fabrication Processes
e.g., fine-grain superplastic forming of complex-shaped articles, isothermal precision die-forging, novel casting techniques, and advanced joining techniques have been developed to make net or near-net shape products, which are very effective in terms of cost/performance trade off. Another new technology that affords promise for broadening the applicability of aluminum alloys is rapid solidification processing, which allows constitutional and microstructural control not possible by the conventional ingot metallurgical route. Yet another, more sophisticated technology is mechanical alloying, which involves processes for mixing elemental powder particles to produce controlled, extremely fine microstructures using a highenergy rolling mill such as an attritor. A further significant improvement in performance over conventional alloys can be achieved by using metal matrix composites (MMCs), in which very high strength and modulus reinforcing fibers or particulates such as SiC, A12O3, boron, and graphite embedded in a high-strength aluminum alloy matrix offer the unique combination of strength, stiffness, low coefficient of thermal expansion, excellent wear resistance, and high temperature stability (Clyne and Withers, 1993). This top, ic is treated further in Sect. 5.4.11. To meet the demand for lighter, higher temperature performance, the aluminides of AlTi3 and AITi are now being investigated because of their combination of low density and elevated temperature capability. 5.3.5.1 Rapid Solidification Processes (RSPs) There are principally two rapid solidification processes (RSPs), i.e., making powder by atomization, and making continuous foil or ribbon by melt spinning.
237
Gas atomization is used to produce smooth, spherical, or rough, irregular powders. The alloy is melted, if necessary in vacuum, and poured through a small diameter refractory nozzle under an applied pressure. Below the nozzle gas jets impinge on the metal stream, breaking it into small droplets which are rapidly cooled by the gas and collect asfinepowder in the bottom vessel (Fig. 5-18). The atomizing gas can be air, nitrogen, argon, helium, or hydrogen. The solidification is only moderately rapid; between 102 and 104 K/s. This process can economically produce a large amount of powder (see also Vol. 15, Chap. 2 of this Series). On the other hand, the melt-spinning process can obtain a rapid solidification rate of between 103 and 105 or 107 K/s in thin continuous foil or ribbon of 30-80 jjxn thickness. The metal stream is poured onto the surface of the single drum or into a gap between twin drums which are watercooled and rotating at speeds of up to 3500 rpm, as shown schematically in Fig. 5-19. There are various RSPs whose cooling rates vary from 103 to 107 K/s, and the products are in the form of continuous foil, ribbon, or wire, or discontinuous
Melt
Atomized powder
Figure 5-18. Inert-gas atomization (Staniek et aL 1980).
238
5 Aluminum-Based Alloys Pressure
Pressure .Melt
- Melt Continuous filament or strip
Water cooled copper rotating drum A) Single drum method
Continuous filament or strip
B) Double drum method
Figure 5-19. Melt-spinning methods (Staniek et al, 1980).
flakes or fibers with a thickness of 10100 jam. The principal features of the microstructures are gradually changing with the higher rates of solidification, and are char-
acterized as follows: (1) microcrystalline structure, completely homogeneous, ultrafine grain size that can be reduced to 0.1 jam or less; (2) refined segregation and very small dendrite-arm spacing; (3) high supersaturation achievable; (4) the creation of nonequilibrium structures and new metastable phases; and (5) the formation of amorphous structures. Figure 5-20 shows the change in the microstructure on increasing the rate of solidification from (A) to (D). The most common methods for consolidating RS powders are hot isostatic pressing and hot extrusion. When superplasticity is applicable, the extruded billets can be deformed into near-final shapes by isothermal forging. The method of consolidation is very important as regards the properties of the final product, because the unique
Figure 5-20. Micro structures of Al-8%Fe alloys solidified using various cooling rates: (A) slowly cooled, (B)-(D) rapidly cooled with increasing rates from (B) to (D).
239
5.3 Advantages in Manufacturing and Fabrication Processes
structure-property relationships resulting from rapid solidification will be in real danger of collapse if inadequate consolidation is obtained and followed by a subsequent heat-treatment. 5.3.5.2 Mechanical Alloying Process (MAP) Starting with a mixture of different powders the aim of the MAP is to produce a new powder. This is done by milling using high-energy ball mills such as attritors (Fig. 5-21) or gravity rolling mills. All particles of the resultant powder have the same chemical composition and microstructure. The MAP was introduced by Benjamin (1970) to produce dispersionstrengthened superalloys. Nowadays it is used worldwide to develop new materials which cannot be obtained by conventional techniques (Arzt and Schultz, 1989). The mechanisms of the MAP are not fully understood. According to Watanabe (1988), the alloying is performed by ball-ball collisions over several stages. In the first stage, individual particles change into flakes by microforging, but cold welding does not occur yet, as shown in Fig. 5-22 a. Then, in the second stage, work-hardened flakes fracture and cold welding occurs to give finer, lamellar composites. Thus repeated kneading results in convoluted, finer lamellae within the particles along with the beginning of dissolution and solid-solution formation, as shown in Fig. 5-22 b. Finally, in the third stage, the lamellae rapidly become finer and more convoluted with a lamellar spacing of less than one micrometer by repeating kneading. The particles then have an extremely deformed, metastable structure which can contain finer dispersoids. With further processing mechanical alloying produces equiaxial particles which contain random
Gas seai
Inert gas
Water jacket
Coolingwater outlet
Rotating wheel rod
Steel or ceramic balls Coolingwater inlet Figure 5-21, Schematic diagram of attritor highenergy ball mill (Staniek et al., 1980).
Pancakeshaped
Flakeshaped
Section
Q (a) 1st stage
Flat
Lamellar structure
(b) 2nd stage Equi-axial particle
Randomized lamellae structure
(c) 3rd stage
Refinement homogenization
Figure 5-22. Schematic diagram of the three stages of mechanical alloying (Watanabe, 1988).
240
5 Aluminum-Based Alloys
lamellae and substructures with dimensions of several tens of micrometers. In the case of aluminum alloys, organic process additives give an additional benefit (Holve et al., 1968), because dispersoids (e.g., A14C3) are formed during the milling process (reactive milling). Carbide dispersoids are intentionally formed in the matrix by the addition of small amounts of carbide-forming elements such as titanium. The presence of such fine dispersoids is of great advantage to MA-processes materials, producing a higher temperature capability than for similar RS-processed materials.
5.4 Main Aluminum Alloy Groups
Table 5-2. Wrought and cast aluminum alloy designations. Wrought alloys
Cast alloys
Designation
Major alloying elements
Designation
Major alloying elements
lxxx 2xxx 3xxx 4xxx 5xxx 6xxx 7xxx 8xxx
None Cu Mn Si Mg Mg and Si Zn Other than above
100 110-199 200-299 300-399 400-499 500-599 600-699 700-799
None Si Cu Mg Zn Mn Ni Sn
9xxx
5.4.1 Alloy and Temper Designations 5.4.1.1 Alloy Designations
5.4.1.2 Temper Designations
Aluminum materials are divided into three types, namely superpurity, commercial purity, and alloys, and are used as either castings or wrought products. Table 5-2 gives the American Aluminum Association (AA) Alloy Designation System for wrought and cast alloys. Each wrought alloy is given a four-digit number of which the first digit is assigned on the basis of the major alloying element. Modifications to the original alloy and impurity limits are indicated by the second digit. The last two digits indicate a minimum aluminum percentage in the case of the lxxx group and serve to further identify aluminum alloys in the 2xxx through 8xxx groups. For an experimental alloy the prefix "X" is added. The designation system for cast aluminum and its alloys has similarities with that for wrought alloys. Again the first digit indicates the major alloy group. The second two digits indicate the different alloys in the group.
A system of temper designations has been adopted in order to specify the mechanical properties and the method by which to achieve them. The system deals separately with either nonheat-treatable, work-hardened alloys or heat-treatable alloys. The main designations are as follows: F: as-fabricated, O: annealed, wrought product only, H: cold-worked, strainhardened, HI: cold-worked only, H2: cold-worked and partially annealed, and H3: cold-worked and stabilized. The second digit represents residual hardening, e.g., the severely cold-worked or fully hard condition is designated HI 8, which is equal to about a 75% reduction in the original cross-sectional area. Thus 2: % hard, 4: Vi hard, 6: 3/4 hard, 8: hard, and 9: extra-hard when used as the second digit. A different system of designations applies for heat-treatable alloys. Tempers other than 0 are represented by the letter T followed by one or more digits. The main designations are as follows:
241
5.4 Main Aluminum Alloy Groups
T: heat-treated, T2: annealed, cast product only, T3: solution plus cold work, T4: solution plus natural aging, T5: artificially aged only, T6: solution plus artificial aging, T7: solution plus stabilizing, T8: solution plus cold work plus artificial aging, T9: solution plus artificial aging plus cold work, and T10: artificial aging plus cold work. 5.4.2 Superpurity and Commercial Purity Aluminum (lxxx)
This group includes superpurity (SP) aluminum (higher than 99.99%) and the various grades of commercial purity (CP) aluminum containing up to 1 % impurities or minor additional elements. The UTS of annealed 99.99% aluminum is about 45 MPa, with a YS of about 10 MPa and an elongation of about 50%. The various grades of CP aluminum have distinctly higher strength afforded by the various strain-hardening tempers. Since iron and silicon are common impurities in CP aluminum, the phase equilibria in the Al-rich corner of the Al-Fe-Si system has been intensively investigated by many authors. Figure 5-23 was produced from investigations by Phillips (1959), where the four phases of Al 3 Fe, oc(AlFeSi), p (AlFeSi), and Si are found to exist in equilibrium with the aluminum sol-
id solution. The solid solubility of Si in aluminum is relatively large, and its maximum solubility is reported to be 1.65 wt.% at 850 K (577 °C), while the solubility of iron in aluminum is very small, as shown in Table 5-3 (Nishio e t a l , 1970). There are several metastable Al-Fe compounds besides the Al 3 Fe stable compound in the Al-Fe binary system, namely Al 6 Fe, Al^Fe, AlmFe, and Al 9 Fe, which are crystallized in the order of increasing solidification rates. The metastable Al 6 Fe and AlmFe that cause the so-called "fir" surface pattern of anodized sheets are known to transform to the stable Al 3 Fe phase on annealing at temperatures higher than 500 °C (773 K), and hence such a pattern can be eliminated (Kosuge and Takada, 1979). Table 5-3. The solid solubility of iron in aluminum (wt.%) (Nishio et al, 1970). Temperature (°Q
Electrical resistivity method
Mossbauer effect method
655 640 630 600 570 550 538 500 450
0.052
_ 0.051 0.043 0.033 0.021 0.017 0.011 0.005
0.025 0.013 0.006 -
Figure 5-23. Liquidus surface of the Al-rich corner of the Al-Fe-Si system (Phillips, 1959). 12
242
5 Aluminum-Based Alloys
5.4.3 Al-Si Alloy System The aluminum-silicon system forms a simple eutectic with a composition of 12.5 wt.% Si at 577°C (850 K) between an aluminum solid solution containing 1.65 wt.% Si and virtually pure silicon. Alloys with silicon as the major alloying addition are the most important of the aluminum casting alloys mainly because of their high fluidity, low shrinkage in casting, high corrosion resistance, good weldability, easy brazing, and low coefficient of thermal expansion (CTE). Moreover, the hardness of silicon particles imparts excellent wear resistance, though machining may present difficulties. Commercial alloys are available with hypoeutectic, eutectic, and hypereutectic compositions. The hypereutectic alloys have recently been more widely employed. Alloys with the coarse eutectic or primary silicon particles exhibit low ductility because of the brittleness of the large silicon plates. For eutectic alloys having the normal microstructure, as shown in Fig. 5-24 a, a refinement or "modification" treatment, discovered by Pacz in 1920, is achieved by the addition of sodium salts or small quantities of 0.005-0.15% metallic sodium or strontium to the melt. As little as 0.001 wt.% of such additives may be enough to modify the microstructure, as illustrated in Fig. 524 b. However, for hypereutectic alloys containing 20 wt.% or more silicon, a sodium addition is ineffective, but phosphorus can be added in the form of phosphorus salts or metallic phosphorus at a level of 0.010.03 % to introduce small, insoluble particles of A1P which serve as nuclei and refine the primary silicon. Figure 5-25 shows the microstructures, (a) without treatment and (b) with treatment by P, for an Al-20 wt.% Si alloy (Adachi; 1984). There has been a great deal of research on the mechanism of
*s Figure 5-24. Optical microstructures of the A l 12 wt.% Si eutectic alloy: (a) unmodified, and (b) modified, x200
such a modification. Induced twinning in the silicon phase is a favorite candidate (Hellawell, 1990). 5.4.4 Al-Mn-(Mg) Alloy System (3xxx) In the binary Al-rich-Mn system, the eutectic temperature, its composition, and the solid solubility are generally confirmed to be 658.5°C (931.5 K), 1.95 wt.% Mn at the eutectic temperature, and 0.36 wt.% Mn at 500 °C (773 K), respectively. The intermetallic compound that is in equilibrium with the aluminum solid solution has a composition closely corresponding to the composition of Al 6 Mn, and crystallizes as a primary phase from the liquid phase with 1.95-4.1 wt.% Mn. This corresponds to the liquid composition of a peritectic reac-
5.4 Main Aluminum Alloy Groups
Figure 5-25. Optical microstructures of the A l 20wt.% Si alloy: (a) without, and (b) with treatment; x200.
tion at 710 °C (983 K), in equilibrium with the intermetallic compound in Al 4 Mn, as shown in Fig. 5-26. The only metastable phase, "G", which has a composition close to Al 12 Mn (14.5 wt.% and 7.69 at.% Mn), is found in rapidly cooled alloys on annealing at 550 °C or lower. The binary Al-Mn alloys containing up to 1.25% Mn are commercially employed
243
as nonheat-treatable alloys designated 3003. A commercially more widely used alloy which has a further addition of about 1.2 wt.% Mg (3004), is particularly used in cans for the beverage industry, as previously mentioned. Since a commercial 3004 alloy usually contains about 0.4-0.5 wt.% Fe and the solid solubility of Fe in Al 6 Mn is relatively high, the 3004 DC-cast ingot contains the Al6(Mn,Fe) compound, which transforms into a-Al 15 (Mn,Fe) 3 Si 2 when Si is also incorporated on annealing at higher temperatures. The presence of finer manganese-containing compound particles results in some dispersion hardening, and thus the UTS of annealed 3003 is typically 110 MPa with corresponding increases with work-hardening tempers. The UTS of a 3004 alloy is increased further to 180 MPa in the annealed condition; there is also an increase in the recrystallization temperature of about 50-60 °C due to an addition of magnesium. For the successful manufacture of 3004 aluminum alloy sheet into beverage cans, control of earing is a critical prerequisite, since ears interfere with the smooth operation of the automated can-making process. To achieve adequate strength in the finished can, sheet in a highly cold-rolled (over 80%) condition has to be used. However, a result of this heavy rolling is that the sheet adopts a cold-rolled temper, which tends to induce ears in drawn cups
Atomic percent maganese 2
800
3
L 4.1
2 700
660.37°
^ . ^
-
—
-
~
•
.^
— L+r
_«——"—•
——
——
710° 659°
1.9
Figure 5-26. The aluminum-rich corner of the binary Al-Mn system (Metals Handbook, 1973).
i
(Al)
1-600
(AD+/9 500
•
Al
/
3 4 5 6 7 Weight percent manganese
10
244
5 Aluminum-Based Alloys
situated +45° to the rolling direction, as shown in Fig. 5-27 a. If the sheet can be processed so that it has a strong cube texture before cold rolling, a tendency for 0/90° earing (Fig. 5-27 b) is induced at this stage which then gradually transforms to 45° earing with increasing cold reduction. Thus a balanced texture giving little or no earing tendency can be achieved in the final sheet product, as shown in Fig. 5-27 c. (See also Vol. 15, Chap. 10, Fig. 10-25 of this Series.) The control of earing in can stock remains a severe practical problem despite much recent research. However, there are some fundamental insights that lead to low earing in the final temper, HI9 condition of 3004 alloys (Nes, 1985). The formation of a strong cube texture is closely related to the recrystallization behavior during ingot homogenization and subsequent hot working. The recrystallization process in 3004 alloys is controlled by the following three factors: the amount of solute in solid solution, the character of the coarse constituent particles such as Al 6 (Mn,Fe), and the size, density, and distribution of fine precipitated dispersoids such as a-Al 15 (Mn,Fe) 3 Si 2 present in the alloy immediately prior to annealing. The above factors further affect the earing behavior through the following processes: (i) homogenization can be carried out at those temperatures and times that enable
(a)
(b)
(c)
Figure 5-27. Deep-drawn aluminum cups showing (a) 45° earing, and (b) 0/90° earing, and (c) no earing with respect to the rolling direction.
the development of a strong cube texture after hot working and annealing; (ii) a well-defined, polygonized dislocation substructure must be developed by an appropriate hot-working procedure to produce a strong cube texture during annealing; and (iii) the low solute supersaturation produced by high annealing temperatures favors the development of the cube texture and subsequently leads to low earing in the H19 sheet. 5.4.5 Al-Mg Alloy System (5xxx) The Al-Mg binary system is the basis for a widely used class of non-heat-treatable aluminum 5xxx series alloys containing from 0.8 to slightly more than 5 wt.% Mg. In aluminum-rich alloys, the simple eutectic reaction occurs at 451 °C (742 K) at a composition of 35.0 wt.% Mg between the aluminum solid solution of 14.9 wt.% Mg and the stoichiometric compounds of Al 2 Mg 3 (37.3 wt.% Mg), at a composition outside its limit of existence (34.837.1 wt.% Mg). Although magnesium has substantial solubility in solid aluminum, the binary alloys do not show appreciable precipitation-hardening characteristics at compositions below 7 wt.% Mg. In the precipitation process, /?-Al8Mg5 (36 wt.% Mg) is observed. Magnesium provides substantial strengthening, and the work-hardening rate in particular increases rapidly with increasing magnesium content. For example, the 5005 alloy (Al-0.8wt.% Mg) has a UTS of 125 MPa, a YS of 40 MPa, and an elongation of more than 25 % in the annealed condition, while the fully workhardened 5456 alloy exhibits 385 MPa UTS, 300 MPa YS, and 5% elongation. It is important to note that these alloys can exhibit some structural instabilities in the following two ways:
245
5.4 Main Aluminum Alloy Groups
(1) Since (}-Al8Mg5 has a tendency to precipitate preferentially in slip bands and at grain boundaries, when the Mg content exceeds 3-4wt.% intergranular attack and stress-corrosion cracking may occur in corrosive environments. In addition, Pphase precipitation occurs slowly at room temperature, and is accelerated if the alloys are in a heavily cold-worked condition or heated slightly above room temperature. Thus small amounts of chromium and manganese are added to most binary alloys in order to increase the recrystallization temperature. These additions also increase the tensile properties by dispersion hardening. Thus a 5054 alloy containing only 2.7 wt.% Mg together with 0.7 wt.% Mn and 0.12 wt.% Cr can have an almost equivalent tensile strength to a binary alloy with as much as 4 wt.% Mg, without any thermal instability. (2) The work-hardened alloys may undergo what is known as "age-softening" at room temperature. The amount of softening increases with increasing cold-working rates of Mg concentrations. This age-softening is generally explained in terms of a relaxation process or preferred precipitation of P-phase on slip bands. A series of stabilized H3 tempers has been adopted to inhibit this effect in practice. During stretching or forming of Al-Mg and some other aluminum alloys, stretcher-strain markings (Luders bands) typically occur in one of two forms, which are associated with the discontinuous or uneven yielding often observed in the stressstrain curves of annealed or heat-treated solid-solution alloys, as illustrated schematically in Fig. 5-28. The strain markings are classified into two types as follows (Thomas, 1966; Pink and Grinberg, 1984): (1) Type A or "flamboyant" markings (as shown in Fig. 5-29 a) occur after a small amount of strain, and are sensitive to
A
t
to
a
/
^|W|Tl1
b' c
CO
1 r\r
^
r*^
B
^
A+ B
Strain, e — * Figure 5-28. Schematic illustration of stress-strain curves associated with various types of discontinuous or uneven yielding (Pink and Grinberg, 1984).
grain size and to the presence of magnesium at the grain boundary. (2) Type B or CB markings also called "parallel band" occur after a large amount of strain and appear as diagonal bands oriented approximately 50° to the tension axis. The bands move up and down the axis during stretching (as shown in Fig. 5-29 b), and terminate at the grip end of the specimen. Failure usually takes place at a Liiders band. The explanation for the parallel bands is frequently given in terms of the Cottrell concept of dislocations moving at a steady velocity and dragging a solute atmosphere along with them at the same velocity. There is a limiting velocity at which the atmosphere can be dragged and, if the dislocation velocity exceeds this, the dislocation will break away and the flow stress will fall until other solute atoms diffuse to reform the atmosphere. This repetitive process provides a simple explanation for the serrated flow curve. The initial smooth part of the curve, which occurs before the jerks begin, has been interpreted as being the amount of strain required to produce a certain excess of vacancies which are necessary to increase the solute
246
5 Aluminum-Based Alloys
Figure 5-29. Stretcherstrain markings: (a) Type A or flamboyant surface markings, (b) Type B or parallel band surface markings (Thomas, 1966; with kind permission from Elsevier Science Ltd.).
atom diffusion to a critical value which enables the solute to keep up with the moving dislocation. Al-Mg alloy sheets are especially susceptible to Luders band formation, which is undesirable since they cause an uneven and rough surface on the products. A slight final plastic deformation such as skin-passing or roller-leveling is performed to keep the dislocations away from the solute atmosphere in order to prevent stretcher-strain markings. Al-Mg alloys have a very low sensitivity to weld-cracking compared with 2xxx, 6xxx, and 7xxx series alloys, and thus they are widely used for welded applications. Structural plates with 5083-0 temper are used, for example, for large welded pressure vessels for carrying liquid natural gas (LNG), where cryogenic storage is involved. They can be polished to a bright surface finish, particularly if made from high-purity aluminum, and thus are used for ornamental articles as well as architectural components.
5.4.6 Al-Cu and Al-Cu-Mg Alloy Systems (2xxx) 5.4.6.1 Al-Cu Alloy System Relatively few commercial alloys based on the binary Al-Cu system are actually used at present, although the sequences of the precipitation process, particularly GP zone formation, have been studied until recently in greater detail for this system than for any other system (Martin, 1968). The decomposition sequence for the supersaturated solid solution (SSSS) of the A l Cu binary system is known to be (SSSS)
GP (1) - GP (2) 0'-CuAL
Tables 5-4 and 5-5 show the nominal compositions and mechanical properties, respectively, of the same commercial alloys. Alloy 2011 containing 5.5% Cu shows good machining characteristics with excellent chip formation owing to the addition of a small amount of the insoluble elements Bi and Pb. Alloy 2025 is used for forgings. Alloy 2219, which contains a larger amount of Cu (e.g., 6.3%), has a relatively high tensile strength at room
247
5.4 Main Aluminum Alloy Groups
Table 5-4. Chemical composition of wrought Al-Cu and Al-Cu-Mg alloys. Designation
2011 2025 2219 2014 2017 2024 2036 2048
Chemical composition (wt.%) Cu
Mg
Mn
Cr
5.5 4.45 6.3 4.4 4.0 4.4 2.6 3.3
0.05 0.02 0.5 0.6 1.5 0.45 1.5
0.8 0.6 0.8 0.7 0.6 0.25 0.4
0.10 0.10 0.10 0.10 0.10
Si
Fe
Ti
Zn
0.4 0.53 0.20 0.8 0.6 0.5 0.5 0.15
0.7 1.0 0.30 0.7 0.7 0.5 0.5 0.20
0.15 0.06 0.15 0.15 0.15 0.15 0.10
0.30
Others 0.4 Bi, 0.4 Pb 0.1 V, 0.17 Zr
0.25 0.25 0.25
Table 5-5. Mechanical properties of wrought Al-Cu and Al-Cu-Mg alloys. Designation and temper 2011-T3 2011-T8 2014-O 2014-T4 2014-T6 2017-O 2017-T4 2024-O 2024-T3 2024-T4 2024-T361 2219-O 2219-T42 2219-T62 2219-T81 2036-T4
UTS
YS
(MPa)
(MPa)
37.7 40.7 18.6 42.6 48.0 18.2 42.6 18.6 48.0 42.1 49.5 17.2 35.8 41.2 47.5 32.0
29.4 30.9 9.8 28.9 41.2 6.9 27.4 7.4 34.3 32.3 39.2 7.4 18.6 28.9 39.2 16.0
temperature and good creep strength at elevated temperatures, as well as high toughness at cryogenic temperatures. 5.4.6.2 Al-Cu-Mg Alloy System The phenomenon of age hardening and the first age-hardenable alloy originated from the discovery by Alfred Wilm in 1906, when he produced "duralumin" of which the nominal composition is 4.5%
Elongation (Sheet)
(Bar) 15 12 18 20 13 22 22
20 18 20 13 18 20 10 10 25
Fatigue strength (MPa) 12.5 12.5 9.0 14.0 12.5 9.0 12.5 9.0 14.0 14.0 12.5
10.5 10.5
E (GPa) 7.2 7.2 7.5 7.5 7.5 7.4 7.4 7.5 7.5 7.5 7.5 7.5 7.5 7.5 7.5
Cu, 0.5-1.0% Mg, and 0.5% Mn. The discovery of duralumin was accidental and dramatic. Wilm was testing the properties of various Al-Cu-Mg alloys in order to develop a stronger aluminum alloy using the principle of quench-hardening, as used for steels. At the weekend his laboratory assistant left the unexamined quenched sample for hardness testing. On the Monday he found unusual hardening in the sample which had been stored for two days
248
5 Aluminum-Based Alloys
at room temperature. This is the process known at present as "natural aging". An addition of Mg to Al-Cu alloys accelerates and intensifies the natural aging. From the investigation of ternary equilibrium, the existence of five phases i.e., (Al), 0-Al2Cu, S-Al2CuMg, Al 6 CuMg 4 , and Al 3 Mg 2 , in the ternary Al-rich solid phase equilibrium diagram has been clarified. The equilibrium phases which contribute to the age hardening, change depending upon the Cu/Mg ratio as follows CuAl2
Al2CuMg
1.5:1
Al 6 CuM g4
The precipitation sequences are divided into two processes as follows (Hardy and Heal, 1954)
Table 5-4 and Table 5-5 show the chemical compositions and mechanical properties of the main commercial Al-Cu-Mg alloys, respectively. A slightly modified alloy of duralumin (2017) is still used, mainly for rivets. Several alloys have been developed and are now widely used for aircraft construction. The alloy 2014 attains higher strength on artificial aging owing to the relatively high Si addition. Another alloy (2024), which contains a higher Mg content of 1.5 % and an Si content reduced to impurity levels undergoes significant hardening on room temperature aging, i.e., T3 or T4. This alloy also shows higher artificial age hardening at around 175°C when coldworked prior to aging. 5.4.7 Al-Mg-Si Alloy System (6xxx)
(1) (SSSS) -> GP(1) -• GP(2) -> -> 0'-CuAl2 -> 0-CuAl2 (2) (SSSS) -> GPB(l) -> GPB(2) -> -> S'-Al2CuMg -> S-Al2CuMg The first process proceeds in the alloy of Cu/Mg = 8, and the first and second processes advance simultaneously in the alloys of 4 < Cu/Mg < 8. In the range of 1.5 < Cu/ Mg<4, the quasi-equilibrium of Al-SAl2CuMg may be established, and thus appreciable age hardening occurs in theses alloys. GPB (1) and GPB (2) correspond to GP (1) and GP (2) containing Mg as well as Cu (Bagaryatskii, 1948). The GPB consisting of Cu and Mg atoms may be formed on {110} matrix planes. The apparent acceleration of this formation may result from either complex interactions between vacancies and the two solute atoms or some preliminary pairing of Cu and Mg atoms. The coherent precipitation of S'-Al2CuMg plays an important role in age hardening at elevated temperatures, such as with 2024T6.
Al-Mg-Si alloys are widely used as medium-strength precipitation-hardening structural alloys, since the (Al)-Mg2Si quasibinary system forms and the solid solubility of the Mg2Si decreases from 1.85 wt.% (1.17% Mg, 0.68 wt.% Si) at the eutectic temperature of 595 °C (868 K) to 0.09 wt.% at 200 °C (473 K) (Phillips, 1946). When no Mn or Cr is included, the iron-rich phases, Fe 3 SiAl 12 , Fe 2 Si 2 Al 9 , or a mixture of the two form depending on the proportions of Mg, Si, and Fe. In the commercial alloys, the Mg and Si are added in the stoichiometric ratio corresponding to Mg2Si (Mg:Si = 1.73:1), or with an excess of Si above that necessary to form Mg2Si, because an excess of Mg causes a considerable decrease in the solid solubility of Mg2Si and hence impairs the precipitation hardening. The precipitation sequence of this alloy system has been elucidated by various techniques, particularly X-ray techniques and transmission electron microscopy, as follows: SSSS -• needle-shaped zone contain-
5.4 Main Aluminum Alloy Groups
ing a high concentration of vacancies -• needle-shaped, internally ordered zone -• rod-shaped p'-Mg2Si intermediate phase -» plate-shaped |3-Mg2Si equilibrium phase. The zones are orientated parallel to the <100> directions of the Al matrix and have coherency strain around the needles. The most used commercial alloys are divided into three classes, the compositions of which are illustrated in Fig. 5-30. The first class (6063) involves alloys with balanced amounts of Mg and Si for Mg2Si up to between 0.8 and 1.2 wt.%, and the second one (6061) contains these elements to give in excess of 1.4 wt.% Mg2Si. The alloys in the third class (6005) contain Si in excess of that necessary to form Mg2Si, and the presence of the excess Si favors precipitation hardening by refining either |3'-Mg2Si precipitates or precipitated Si particles. The 6063 alloys can easily by extruded and afford a potential economic saving in the "press-quenching" process (Horiuchi and Kaneko, 1976), in which the product may be quenched at the extrusion press when it emerges from the die, and thus the need for solution treatment can be eliminated. Press-quenching is normally performed by means of water spraying or by a pressurized air blast, and then subsequent aging at 160-190 °C is applied to obtain a moderate tensile strength. With the T6 temper, the typical tensile properties are a YS of 215 MPa and a UTS of 245 MPa. A typical 6061 alloy (Al-1.0 wt.% M g 0.6 wt.% Si-0.25 wt.% Cu) in the second class contains in excess of 1.6 wt.% Mg2Si, and thus it is usually necessary to execute solution treatment and water-quenching as separate operations after extrusion. Because its tensile properties are higher than those of 6036, this alloy is used for generalpurpose structural materials.
0.2
0.4
0.6 0.8 1.0
249
1.2 1.4
Si {%) Figure 5-30. The composition ranges of typical AlMg-Si alloys.
The third class, typical alloy 6005, contains Si in excess of that corresponding to Mg2Si, and this excess Si can enhance the tensile properties, but may reduce the ductility and cause intergranular embrittlement, which is attributed partly to the tendency for Si to segregate at grain boundaries. An addition of small amounts of Mn or Cr may prevent this harmful effect by grain refinement as well as by retarding recrystallization. In Al-Mg 2 Si alloys containing more than 0.9 wt.% Mg2Si, storage at room temperature before artificial aging results in tensile properties inferior to those of the alloy aged immediately after quenching, whereas storage at room temperature has a favorable effect on Al-MgZn 2 alloys. Figure 5-31 shows TEM micrographs illustrating the effect of delayed artificial aging on an Al-1.22 wt.% Mg2Si alloy after aging at 160°C for 24 h: (a) without any delay, and (b) with storage for 3 days at 20°C before aging at 160°C. Figure 5-31 b shows a coarser precipitate structure, which results from natural aging during storage and can lead to a failure to develop the
250
5 Aluminum-Based Alloys
Figure 5-31. TEM micrographs showing the effect of delayed artificial aging on an Al-1.2wt.%Mg 2 Si alloy: (a) after aging at 160°C for 24 h without any delay, and (b) with storage for 3 days at 20 °C before aging at 160 °C (Pashley et al., 1966).
optimum mechanical properties upon subsequent artificial aging. Pashley et al. (1966) explained semiquantitatively the whole effect of the twostep aging process in terms of the stability of clusters formed during preaging at room temperature. As clustering takes place during storage, the supersaturation of solute remaining in solution is diminished. When the supersaturation with respect to both vacancies and solute is further reduced as a result of increasing the temperature to the artificial aging temperature, the size at which a cluster becomes stable is increased, and some clusters will continue to grow whilst others dissolve. Thus the longer the alloy is stored at room temperature, the lower will be the supersaturation at the start of artificial aging and the smaller will be the number of clusters. This is the major factor leading to the formation of a coarser precipitate structure and hence poorer tensile properties as a result of room temperature storage. An addition of copper of ~0.24 wt.% to the alloy considerably reduces the effect of storage, because it reduces the rate of natural aging. 5.4.8 Al-Zn-Mg-(Cu) Alloy System (7000) 5.4.8.1 The Precipitation Sequence The precipitation sequence of the ternary Al-Zn-Mg system has been elucidated by various techniques as (Schmalzried and
Gerold, 1958) SSSS -> GP zones -> r|'-MgZn 2 -> -> r|-MgZn 2 (T-Al 2 Zn 3 Mg 3 ) The GP zones formed near room temperature in Al-2.7 at.% Zn-4at.% Mg have an ordered CuAul structure which is made up of alternate layers of Zn-rich (200) and Mg-rich (200), while at higher temperatures, i.e., 60-150°C, spherical zones are formed in coherence with the (111) plane of the Al matrix. The formation of r|'MgZn 2 in the shape of hexagonal platelets occurs on the (111) plane; and this phase has the orientation relation of {0001 j ^ / / {lll} a and{1010} n 7/{110} a . With regard to decomposition of the quaternary system containing Cu, two different opinions have been advanced. One opinion is that Cu atoms only go into GP zones and r|'-phases rather than generating any essential changes in the decomposition process. The other proposal is that the decomposition process of the Al-Cu-Mg ternary system SSSS -> GPB zones -> S'(Al2CuMg) -> S proceeds in the same way as that of the Al-Zn-Mg ternary system. However, in the case of practical aging treatments only if and v[ have been detected in the A l - Z n Mg-Cu system, with the exception of the experimental result whereby the S' phase was observed at aging temperatures above 175°C.
5.4 Main Aluminum Alloy Groups
251
Figure 5-32. TEM micrographs of the Al-5-9% Zn-2.9%Mg alloy acetone-quenched to — 95 °C, and (a) aged for 3 h at 180°C, or (b) aged for 2 h at 180°C, held, at 20 °C for 192 h and reaged for 2 h at 180°C (Lorimer and Nicholson, 1966; with kind permission from Elsevier Science Ltd.).
5.4,8.2 Two-Step Aging In contrast to the Al-Mg-Si alloys, storage at room temperature in this case leads to the formation of finer precipitate structures at the artificial aging temperature. Lorimer and Nicholson (1966) deduced from TEM experimental results that the aging behavior of Al-Zn-Mg alloys can be divided into three classes which can be defined by the temperature ranges involved: (i) Alloys quenched and aged above the GP zone solvus (i.e., above 155°C for the Al-5.9 wt.% Zn-2.9 wt.% Mg alloy); since no GP zones are ever found during heattreatment, there are no easy nuclei for subsequent precipitation and a very coarse dispersion of precipitates results, with nucleation principally on dislocations. (ii) Alloys quenched and aged below the GP zone solvus (e.g., below 155°C for this alloy); GP zones form continuously and grow to a size where they are able to transform into precipitates. This transformation occurs rather more slowly in the grain-boundary regions owing to the lower vacancy concentration there, but since aging will always be carried out below the GP zone solvus, no PFZ is formed other than a very small [ ~ 300 A (30 nm)] solutedenuded zone as a result of precipitation at the grain boundary. (iii) Alloys quenched to below the GP zone solvus and aged above it (e.g., this
alloy quenched to room temperature and aged at 180°C); this is the most common situation, and the final dispersion of precipitates and the PFZ width are controlled by the nucleation treatment below 155°C, where the GP zone size distribution is determined. A long nucleation treatment gives a fine dispersion of precipitates and a narrow PFZ. The following experimental results can be advanced to account for the microstructures produced by two-step aging. If the alloy is aged at 180°C following a short holding treatment designed to give a wide PFZ, as shown in Fig. 5-32 a, but is then given a further "nucleation treatment" by holding it below the temperature of the GP zone solvus, subsequent reaging at 180°C results in fine precipitation inside the original zone, as shown in Fig. 5-32 b. This is simply explained by Embury and Nicholson (1965): the zones formed inside the PFZ after the initial holding treatment were too small to act as nuclei on subsequent aging at 180°C and hence dissolved; however, the solute atoms can reform as zones after a further low temperature treatment and if the zones are allowed to grow to a sufficient size, they may act as nuclei for precipitation during the second aging period at 180°C. In Al-Zn-Mg-Cu alloys, e.g., 7000 system alloys, aging at higher temperatures of 160-170°C, at which the if phase precipitates, results in a significant in-
252
5 Aluminum-Based Alloys
crease in the resistance to SCC but the tensile properties are much reduced. Subsequently, a two-step aging treatment, designated the T73 temper, is used, by which a finer dispersion of if -precipitates can be obtained through nucleation from preexisting GP zones, as described above. 5.4.8.3 Thermomechanical Treatment (TMT)
TMTs involving a combination of plastic deformation and heat-treatment have been employed to improve the ductility, toughness, and SCC resistance of 7000 series alloys, especially in the short transverse direction, without decreasing the strength in comparison with conventionally processed materials. There are two kinds of TMT - intermediate thermomechanical treatment (ITMT) and final thermomechanical treatment (FTMT).
Inferior transverse properties in conventionally processed materials are attributable to the grain size and elongated shape as well as to the presence of second phase particles at the grain boundaries. ITMTs are aimed at minimizing the influence of the original cast structure by intermediate recrystallization during the working schedule, as shown in Fig. 5-33. FTMT, which is aimed at the establishment of a condition in which the dislocations introduced by plastic deformation interact most favorably with the age-hardening process, can produce materials that have improved properties compared to conventionally aged materials. Intermediate Thermomechanical Treatment (ITMT) ITMTs are divided into two classes; the ISML (Istituto Sperimentale dei Metalli
after full homogenization/ I
"V as cast ingot
r/~\
after conventional hot
conventional
treatment ITMT cycle
after partial homogenization
Figure 5-33. Schematic diagram showing the microstructural transformations of 7075 ingots processed by conventional treatment (top) and by an ITMT cycle (bottom) (DiRusso et al., 1974).
253
5.4 Main Aluminum Alloy Groups
Table 5-6. Comparison of the tensile properties of conventionally and ITMT-processed 7075 plate [lin (2.5 x 10" 2 m) thick] in the L direction (LT direction). Treatment
YS (0.2%)
UTS
(MPa) (a) T6 Temper (1) Conventional (2) FA-ITMT (3) ISML-ITMT
RA
K
(MPa)
£ [2 in (5xlO~ 2 m)] (%)
(%)
(MPa m1/2)
526 (502) 507 (509) 514 (508)
587 (568) 574 (572) 576 (573)
10 (9.5) 18 (19.0) 17.5 (18.2)
14-17(14-16) 29.8 (35.1) 29.4 (29.6)
28.1 (22.6) 30.9 (27.9) 30.4 (33.8)
(b) T73 Temper (1) Conventional (2) FA-ITMT (3) ISML-ITMT
457 (445) 468 (458) 454 (450)
528 (516) 528 (518) 520 (513)
12 (10.5) 16.5 (16.0) 16.5 (14.5)
29 (20.0) 48.5 (45.1) 50.0 (38.4)
34.7 (31.0) 51.3a(44.4)a 51.6a (43.5)a
(c) FTMT treatment (1) FA-ITMT (2) ISML-ITMT
574 (561) 608 (573)
607 (603) 630 (614)
13.7 (12.2) 13.2(11.2)
37.4 (34.6) 28.2 (25.2)
24.9 (22.9) 27.9 (22.6)
KQ (ASTM E-399).
Leggeri)-ITMT (DiRusso et al., 1974) is designed mainly for thin sheet, i.e., 4 mm thick, while the FA (Frankford Arsenal)ITMT (Waldman et al., 1974) is mainly for thick plate, i.e., 25 mm thick plate. (i) ISML-ITMT: (a) Partially homogenized and worked 70-80 % at a relatively low temperature, i.e., 260-330 °C, maintaining most of the Cr in supersaturated solid solution; (b) recrystallization while preventing grain-boundary migration by a fine dispersion of E-(Al 18 Mg 3 Cr 2 ) precipitated from supersaturated solid solution; (c) homogenization; and (d) conventional hot working to obtain a fine, equiaxed grain structure. (ii) FA-ITMT: (a) Complete homogenization of precipitate Cr as E-phase particles; (b) slow cooling to precipitate Zn, Mg, and Cu as coarse particles; and (c) working at a low temperature and homogenization to obtain a fine, equiaxed grain structure.
Final Thermomechanical Treatment (FTMT) The optimum FTMT conditions for 7075 alloys are as follows: (a) Solution treatment at 465-470 °C; (b) quenching in room-temperature water; (c) natural aging for 2-3 days; (d) first artificial aging at 105 °C for 6 h; (e) plastic deformation; and (f) final artificial aging at 105-120 °C for various times, depending upon the degree of prior deformation. It has also been shown that the satisfactory hardening response brought about by the sequence mentioned above is due to the cooperative action of dislocation substructures and to a homogeneous, dense distribution of fine x\r precipitate particles. Tables 5-6 compares the tensile properties of conventionally and ITMT-processed 7075 plate (1 inch thick) in the L (longitudinal) and LT (long transverse) directions.
254
5 Aluminum-Based Alloys
5.4.8.4 Commercial Al-Zn-Mg-(Cu) Alloys This alloy system is divided into two categories. One contains medium-strength alloys with reduced Zn and Mg contents and little or no Cu addition, with the advantage of being readily weldable. The other contains high-strength alloys containing the quaternary addition of Cu as well as higher Zn and Mg contents, as shown in Table 5-7. (i) Weldable Al-Zn-Mg alloys can ageharden significantly at room temperature, and moreover are relatively insensitive to the rate of cooling from a high temperature. These characteristics are very suitable for the welding process, and thus there is considerable strength recovery after welding, without further heat-treatment. It is generally accepted that the Zn-Mg content should be less than 6 % for a weldable alloy to have a satisfactory resistance to SCC, though the tensile strength is known to increase as the Zn + Mg content is raised. An addition of smaller amounts (0.1-0.3%) of one or more of the transition elements such as Cr, Mn, and Zr is also necessary to improve the SCC resistance. These elements can help to control the grain structures during fabrication and heat-treatment, while a Zr addition can also improve the weldability. The amount of Cu has to be kept below 0.3 % in order to minimize hot-cracking during welding, as well as corrosion in service. With regard to heat-treatment, slower quenching rates such as air cooling from the solution-treatment temperature minimize the residual stresses and decrease the difference in the electrode potentials throughout the microstructure, and hence can lead to improved resistance to SCC. Another method is to age the alloy artificially, e.g., the T73 temper, as described previously. The
compositions of representative alloys, e.g., 7003 and 7N01, and their mechanical properties are shown in Tables 5-7 and 5-8, respectively. (ii) High-strength A l - Z n - M g - C u alloys have received special attenti