Structure and Properties of Nonferrous Alloys

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.

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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.

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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-

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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

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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

]/c)

EZ33./;.

40-

ZE41-

f-WE43 1

30-

Li

I

c CO

AE42v

/EQ21

/

f^

_

T

10-

Xw

E54

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

'.

'.

iiiiiT,

;

:

i , ,

147

Mii

b)

Room Temperature

120

120

WE43 /WE54 too

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

j I

! I

:

b) 250'C 150

>.^- > ... : .. r r > > ^ s : ^ ^

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^"*-

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—

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to QE22^ > "1^^CZ O

Q. CD CD

O

^EZ33 50

50 •AE42

;

0 1.000

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

a)

•

- - WE43

b)

— AE42 H

400-

400-

Mg-Sm6

• Mg-Gd9.3Y4.8Mn0.6

•

•

g>300-

300-

en

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200-

2000)

03 N \

100

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s

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v

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100-

Mg-Sm6

• Mg-9.3Gd4.8Y0.6Mn 1

100

200

300

400

I

I

I

100

200

300

400

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

x x x x x x x x

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a

© o o o o o o o

o

CO

CO

CO

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CO

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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).

X ^

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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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151

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. ^

200

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

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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.

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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.

224

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

232

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).

234

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 attention because they have the greatest response of all the aluminum alloys to age hardening. Because of a high susceptibility to SCC, this problem has been the subject of continuing research, as described in Sect. 5.2.3.5. Research has been carried out into improving the resistance to SCC with attention to both alloy composition and aging treatments. The 7075 composition is the best known, while, e.g., 7178-T6 is used for compressively stressed members and 7079T6 is suitable for large forgings. A single aging treatment, T6 temper (aging at a temperature between 120 and 135 °C), which precipitates GP zones, results in a high response to hardening and hence incurs a relatively high susceptibility to SCC. Subsequently, a two-step aging treatment designated the T73 temper is commonly used to prevent SCC, while to increase the resistance of 7000 alloys to exfoliation corrosion, another two-step aging designated the T76 temper (4 h/ 121 °C +18 h/163 °C) is now applied. With regard to RRA heat-treatment, the details are explained in Sect. 5.2.3.5. The composition of representative high-strength alloys and their mechanical properties are shown in Tables 5-7 and 5-8, respectively.

5.4.9 Al-Li Alloy System Aluminum alloys containing lithium additions are considered potential structural materials, especially for aerospace applications. Each weight percent of Li added to an aluminum alloy reduces the density by approximately 3 % and increases the elas-

255

5.4 Main Aluminum Alloy Groups

Table 5-7. Chemical composition of wrought Al-Zn-Mg-(Cu) alloys. Designation

7003 7010 7050 7075 7175 7475 7079 7N01

Chemical composition (wt.%) Zn

Mg

Cu

Si

Fe

Cr

Mn

Zr

5.8 6.2 6.2 5.6 5.6 5.7 4.3 4.3

0.8 2.5 2.3 2.5 2.5 2.3 3.3 1.5

0.20 1.8 2.3 1.6 1.6 1.6 0.6 0.2

0.30 0.10 0.12 0.40 0.15 0.40 0.30 0.30

0.35 0.15 0.15 0.50 0.20 0.50 0.40 0.35

0.20 0.05 0.04 0.20 0.20 0.22 0.20 0.30

0.30 0.30 0.10 0.30 0.10 0.06 0.20 0.45

0.15 0.14 0.12 0.25

Ti 0.20 Ni 0.05 0.06 Zr): 0.25 0.10 0.06 0.10 0.20 V0.10

Al balance balance balance balance balance balance balance balance

Table 5-8. Mechanical properties of wrought Al-Zn-Mg-(Cu) alloys. Designation

7003 7010 7050 7075 7075 7075 7175 7475 7475 7475 7079 7079 7N01 7N01 7N01

Temper

T5 T73651 T73651 O T6, T651 T73, T73651 T736 T6, T651 T76, T7651 T73, T7351 O T6, T651 T4 T5 T6

Fatigue strength (MPa)

E

(MPa)

Elongation (%)

26.5 45.9 46.9 10.7 52.6 45.4 47.4 51.5 48.0 45.4 10.7 49.0 23.0 30.6 30.6

15 11 13 17 11 13 11 12 12 14 17 14 16 15 15

13.0 16.3 16.3 16.3 21.9 16.3 13.2 13.2

7.4 7.4 7.3 7.3 7.4 7.4 7.4 7.4 7.3 7.3 7.3 7.4 7.4 7.4 7.4

UTS

YS

(MPa) 32.7 53.1 53.1 23.0 59.7 52.6 54.6 57.7 54.6 52.6 23.5 56.1 37.2 35.7 37.8

tic modulus by approximately 6 % for Li additions up to 4 wt.%. In the 1950s researchers at Alcoa developed the high-strength Al-Li-Cu alloy 2020 (Al-1.1 Li-4.5 Cu-0.2 Cd-0.5 Mn). However, this alloy had low ductility and fracture toughness when using the maximum strength temper. These limitations and production problems led to its withdrawal as a commercial alloy in 1969 (Balmuth and Schmidt, 1980).

(GPa)

Since 1973, the rapid increase in fuel costs has accelerated the research into developing more fuel-efficient aircraft and hence on developing advanced Al-Li alloys which can reduce the aircraft weight. On the other hand, owing to the progress in fundamental research on fracture mechanics, a good prospect for the improvement of the ductility of Al-Li alloys has led to the control of the microstructures of the alloys. Extensive research has resulted

256

5 Aluminum-Based Alloys

very complicated one and is not yet completely known. A variety of age-hardening precipitates, including 5'-Al3Li and 5Al3Li in the Al-Li binary system and S'(Al2CuMg) and T1(Al2CuLi), as well as the metastable or stable oc'-Al3Zr phase, are possible.

in the development of three representative Al-Li alloys which are registered with Aluminum Association Designations; the chemical composition and mechanical properties of these alloys are given in Table 5-9 and Table 5-10, respectively. 5.4.9.1 The Precipitation Sequence

5.4.9.2 Precipitate Structure and Fracture Toughness

Precipitation in the Al-Li-Cu system, e.g., 2020 and 2090 alloys, follows the following scheme (Tamura etal., 1970; Suzuki etal., 1981, 1982) 5'-Al3Li GP(1) -> GP(2) -+ 0'-CuAl2 - 9-CuAl2 Ti(Al 2 CuLi) -» T1(M2CuLi)

ssss

The precipitate phases vary depending on the ratios of Li and Cu. The copper precipitates independently from the Li and follows the sequence that occurs in an Al-Cu binary system. Simultaneously, Li precipitates as 5'-Al3Li and Ti(Al 2 CuLi), as well as 5-Al3Li and T±. The precipitation sequence of the Al-LiCu-Mg-(Zr) system, e.g., alloy 8090, is a Table 5-9. Chemical composition of Al-Li -Cu-(Mg) alloys. Chemical composition (wt.%)

Designation Li 2090 8090 8091

Cu

Mg

Zr

2.25 2.75 <0.25 0.12 1.10 0.12 2.40 1.35 0.85 0.12 2.60 1.90

Fe

Si

<0.12 <0.10 <0.15 <0.10 <0.30 <0.20

Additions of small amounts of Zr to A l Li-Cu and Al-Li-Cu-Mg alloys and a suitable thermomechanical treatment result in beneficial micro structures for obtaining increased fracture toughness. Zr precipitates coherently and homogeneously as a'-Al3Zr, which controls the grain size and shape. Moreover, oc'-Al3Zr is known to serve as a nucleation site for 8'-Al3Li, and hence the 5'-Al3Li forms a complete envelope around the oc'-Al3Zr precipitated during homogenization and hot working, and leads to the formation of composite precipitates. The composite precipitates cannot be sheared by moving dislocations, and this results in dispersed slip bands and consequently high ductility and fracture toughness (Gayle and Vander Sande, 1984). In commercial alloys containing Cu there are two kinds of precipitate, 5'-Al3Li and T1(Al2CuLi). Figure 5-34 demonstrates the dramatic change in the microstructure of an Al-Li-Cu alloy after several percent stretching before aging

Table 5-10. Mechanical properties of Al-Li-Cu-(Mg) alloys.

Designation

Temper

UTS (MPa)

YS (MPa)

Elongation (%)

K (GPa/m2)

E (GPa)

Density (g/cm3)

Alloys being substituted

2090 8090 8091

T8 T8 T8

569 476 560

530 400 520

7.9 9.0 4.0

42.5 45.6 28.0

78.6 78.6

2.59 2.55 2.55

7075-T6X 2024-T3X 7075-T6X

5.4 Main Aluminum Alloy Groups

257

A good, recent overview of Al-Li-X metallurgy is available by Grimes (1990). 5.4.10 Aluminum Powder Metallurgy Products

Figure 5-34. Effect of stretching on the distribution of Tt phase in precipitate structures of Al-Li-Cu alloys; (a) without stretching, and (b) with stretching (Lee and Frazier, 1988; with kind permission from Elsevier Science Ltd.).

Powder metallurgy (P/M) is the production, processing, and consolidation of fine particles to produce a solid material. Figure 5-35 shows a schematic representation of the powder metallurgy processes for aluminum alloys (Staniek et al., 1981). (See also Vol. 15, Chap. 4 of this Series.) For aluminum rapid solidification (RS) and mechanical alloying (MA) processes (Sects. 5.3.5.1 and 5.3.5.2) produce improved mechanical and corrosion properties compared with those of conventional ingot wrought materials. The refined microstructures obtained by P/M result in improvements in many of the mechanical properties, i.e., room and elevated temperature strength, toughness, fatigue crack propagation resistance, and SCC resistance. 5.4.10.1 Rapid Solidification

(Lee and Frazier, 1988). The dislocations introduced into the matrix by 10 % stretching serve as nucleation sites for the T± phase, and cause a very fine, homogeneous distribution of the phase within the grains (Fig. 5-34 b), while coarse T\ precipitates mainly on the subboundaries in the case of no stretching (Fig. 5-34a). The stretching treatment plays a dominant role in considerably improving the strength as well as the toughness of commercial Al-Li alloys. Recently, small amounts of Sn, Cd, and In (0.14-0.18%) have been reported to have a favorable effect by producing a fine, homogeneous dispersion of Tx phase in A l Cu-Li alloys (Blackburn and Starke, 1987).

Rapid solidification produces improved microstructures by suppressing equilibrium cooling reactions, leading to a final product that can have a finer grain size, a smaller constituent particle size, increased alloy strengthening, and a novel dispersion of fine intermetallic particles. Particles from the RS process have to be degassed and consolidated by various methods (Savage and Froes, 1984). The degassing of a powder at an elevated temperature, usually in a vacuum, evacuates H 2 O and H 2 by evaporating physisorbed H 2 O and decomposing hydrides from the surface oxide layer. At low temperatures of up to 300 °C, either evaporated or decomposed water vapor is the pre-

258

5 Aluminum-Based Alloys Raw material

Melting in vacuum

Atomizing

Heat treatment

Removal of capsule

dominant species outgassed. At higher temperatures, water vapor reacts with Al and Mg to create additional MgO and A12O3, so generating H 2 . In practice the hydrogen content of a 7091 alloy powder compact can be reduced to ~ 1 ppm after degassing at 520 °C. After degassing, the compacts are hot pressed in a vacuum to the maximum density attainable. Hot-pressed billets have to be hot worked, because the surface oxide film is tenacious and cannot be broken suf-

Figure 5-35. Schematic representation of powder metallurgy processes for aluminum alloys (Staniek et al, 1980).

ficiently by isostatic pressing techniques, such as HIP (hot isostatic pressing). 5.4.10.2 Powder Metallurgy (P/M) Alloys Developed by Rapid Solidification (RS) Alloy development has been proceeding in the following three main classes: (i) high room-temperature strength with high corrosion resistance, (ii) low density and high modulus, and (iii) improved high-temperature strength (Froes and Pickens, 1984).

5.4 Main Aluminum Alloy Groups

(i) 7091 and 7090 are A l - Z n - M g - C u alloys which are similar in composition to ingot metallurgy (I/M) alloy 7175, but contain 0.4% Co and 1.45% Co (7090: higher Zn content also), respectively. The cobalt forms Co 2 Al 9 or (Co, Fe) 2 Al 9 particles which are finely dispersed and serve to refine the grain size, and hence these particles enhance the resistance to SCC; the mechanism of which is described in Sect. 5.2.3.5. The alloy 7090-T7E71, for example, may have room temperature longitudinal properties of UTS 637 MPa and YS 595 MPa with 10% elongation. (ii) The alloys with compositions based on A l - L i - C u - M g - Z r are now considered to be the most attractive, although no Al-Li-based alloy has been fully commercialized. One early experimental SiC whisker reinforced Al-Li-based alloy showed a modulus of 170 GPa and a UTS of 1170 MPa, which are over twice the strength and modulus of high-strength 1/ M alloys. (iii) Several RS alloys based on the A l Fe system, e.g., Al-8 Fe-4Ce (Alcoa) and A l - 8 F e - 2 M o (Pratt & Whitney) have been reported to have excellent elevatedtemperature strengths. 5.4.10.3 Mechanically Alloyed Aluminum

In the case of aluminum, carbon derived from process control agents (PCA) is incorporated into the processed powder and reacts with the aluminum to form very fine carbides. These carbides and fine oxide particles derived from the break-up of surface films on the initial powder particles create a dispersion which stabilizes a submicron grain size. Therefore a portion or all of the strengthening may be obtained from the ultrafine grain size stabilized by the oxide and carbide dispersion, rather than relying on precipitation reactions or

259

cold working alone, as in conventionally manufactured alloys. The Inco mechanically alloyed powder (MAP) AL-9052, with the composition A l - 4 % Mg-0.4% O - l . l % C, for example, has been shown to have room-temperature mechanical properties of YS: 380 MPa, UTS: 450 MPa, E\ 76 GPa, and Klc: 44 MPa y/m with 13% elongation. This alloy is produced for applications in marine engineering and shows good corrosion resistance with a weight loss of 0.009 mm/ year in a marine atmosphere and 0.01 mm/ year with alternate immersion in 3.5 % NaCl; it also has an SCC threshold of 380 MPa. Another example is the Al12.5% Ti MA alloy, which has a typical stable, ultrafine grain structure containing fine Al3Ti particles of about 20-250 nm in diameter. This alloy has an E value of 104 MPa - 50% greater than that of conventional aluminum alloys - accompanied by a density increase of 4 % with good ductility and high strength at elevated temperatures, i.e., 220 MPa at 220 °C and 150 MPa at450°C. 5.4.11 Aluminum Matrix Composites 5.4.11.1 Maturity and Intrinsic Characteristics of Aluminum Metal Matrix Composites (MMCs)

Composite materials are defined as the combination of two or more chemically different monolithic materials in order to produce a third material which has unique properties when compared to the monolithic materials' properties. The major advantage of this is the ability to tailor mechanical and physical properties to meet specific needs, which could not be realized by the monolithic materials. For instance, assuming the effect of the interface between the filament and matrix is minor, the

260

5 Aluminum-Based Alloys

Young's modulus of a composite (Ec) consisting of multiple filaments uniaxially aligned will be given by the "rule of mixtures (ROM)".

where Ef and Em are the Young's moduli of fiber and matrix, and V{ denotes the fiber volume fraction. The ROM concept for expressing the strength of both continuous and discontinuous fiber composites has been developed by various investigators, and for uniaxially aligned continuous fiber composites, the following simple ROM is often used as a first approximation. <jc = of Vf +
(5-6)

where oc and of are the ultimate tensile strength of the composite and fiber, and a'm is the matrix stress at the ultimate tensile strain of the fiber. However, the above ROM expression has to be modified somewhat in reality. The composites are classified into three types, i.e., PMCs (polymer matrix composites), MMCs (metal matrix composites), and CMCs (ceramic matrix composites) (Guruganus et a l , 1988). (See also Vol. 13, Chaps. 3,4, and 5 of this Series, respectively. All aspects of MMCs are comprehensively covered in a recent monograph by Clyne and Withers, 1993.) Each composite has its own intrinsic characteristics. Thus the main problem in expending their applications has been the attainment of a performance and cost level that would allow effective utilization in industry of each individual composite. PMCs are the most mature materials because they were first developed in the 1960s, and they have good structural properties, ease of fabrication, and a cost relatively lower than for the other two composite types. The primary advantages of PMCs include good in-plane stiffness and

strength together with a lower density, a low coefficient of thermal expansion, and low electrical and thermal conductivity in addition to the above-mentioned advantages. The major disadvantages are a lower high-temperature strength, moisture sensitivity, and outgassing in space. CMCs are the most recently developed composite system. Their primary advantages comprise high stiffness and strength even at higher temperatures, low density, and good wear resistance. In addition to these, when compared to monolithic ceramics, CMCs have improved fracture toughness, and impact as well as thermal resistance. MMCs lie between PMCs and CMCs in development and maturity. The technologies to produce MMCs depend primarily on the type of reinforcements and the matrix alloys chosen. The present stage of development of each technology depends on the process itself, but there are few MMC forms sufficiently mature to be considered for aerospace, automotive, and leisure goods. The advantages of MMCs include high specific strength, specific stiffness, improved elevated temperature properties, better fatigue resistance, improved wear resistance, high damping properties, and a low coefficient of thermal expansion. In addition, the major advantage of MMCs like other composites is the ability to tailor their specific mechanical and physical properties to meet the needs of products. 5.4.11.2 The Processing and Characteristics of Various Metal Matrix Composites (MMCs) MMCs include a very broad class of materials which may be divided into three main groups: continuous fiber, whisker or short fiber, and particulate-reinforced metals.

5.4 Main Aluminum Alloy Groups

261

Table 5-11. Typical properties of reinforcements. Whiskers: Producer

Type

(3-SiC a-Si3N4 Graphite K 2 • 6TiO 2

Tokai Carbon Tateho Chemicals Nikkiso Otsuka Chemicals

Diameter

Length

Gravity

(urn)

(urn)

(g/cnr 5)

0.1-1.0 0.1-0.6 0.1-1.0 0.2-0.5

50-200 20-200 10-200 10-20

3.19 3.18 2.25 <3.3

Tensile strength (GPa)

Young's modulus (GPa)

3-14 14 7-21 >7

400-700 380 700-800 >270

Young's modulus (GPa)

Short fibers: Type

Producer

Alumina

Nichiasu

Alumina/silica

Isolite

Zirconia

Shinagawa

Chemical composition A12O3 95% SiO2 5% A12O3 47.3% SiO2 52.3% ZrO 2 95% CaO 4%

Diameter

Gravity

(Mm)

(g/cnr

Tensile strength (GPa)

3

3.6

1.1

2.8

2.6

1.3

5

5.8

1.5

120

Long fibers: Producer

Type

filaments

(urn)

(g/cm 3)

Tensile Young's strength modulus (GPa) (GPa)

1 1

100 140 17

2.57 3.0 3.25

3.4 3.4 1.8

Number

Diameter Density

of CYD fiber Multifilaments

B/W SiC/C A12O3 SiC C

AVCO (USA) AVCO (USA) A12O3 85% Sumitomo SiO2 15% Chemical SiC Nippon Carbon PAN Toray Toray Mitsubishi Pitch Chemical

Reinforcement materials for MMCs can be classified into four major groups: whiskers, short fibers, long fibers, and particulates. Reinforcements are generally ceramics (oxides, carbides, and nitrides) which have an excellent combination of strength- and stiffness-to-weight ratios at both room and elevated temperatures be-

1000

500 6000 3000 1000-12000 1000-12000

10-15 7 6.5

2.55 1.76 1.81 2.0 2.1

2.4-2.9 3.5 2.7 1.9 >2.9

390 420 210 180-200 230 390 200 690

cause of their strong interatomic bonding. Typical properties of whiskers and short and long fibers are shown in Table 5-11. The most popular matrix is aluminum, although titanium and magnesium and, to a lesser extent, superalloys and intermetallic compounds are being adopted because their unique mechanical and physical char-

262

5 Aluminum-Based Alloys

acteristics are extremely attractive for various structural as well as nonstructural applications. The manufacturing processes for forming the component materials into the actual composites are known to take many forms. These special manufacturing methods may have a substantial effect on the properties and cost of the composites. Serious challenges have been overcome to seek the appropriate manufacturing method

which will be able to produce a composite with acceptable properties at the lowest cost. There are five main manufacturing processes that have reached industrial status successfully: (a) liquid-metal infiltration, (b) diffusion bonding, and diffusion bonding of plasma-sprayed monolayer tapes, as shown in Fig. 5-36, (c) squeeze casting, and (d) a powder metallurgy method, as shown in Fig. 5-37 (Guruganus etal., 1988).

Precursor wire Pultrusion die Tow

Spool of Tow

Molten-metal bath (a) Liquid-metal infiltration

Wrap fibers on aluminium foil

Lay up desired layers

Encapsulate and in* draw vacuum

and cut to shape Heat to consolidation HI* temperature

Cool, remove, and clean part

Apply pressure and hold for consolidation cycle

nnnnnnnnp

(b) Diffusion bonding Plasma torch ,-, Jl/CompositeXL

Vacuum system

Plasma-sprayed material

11/

Carrier foil

llf mandrel Powder feed - 0 0 0 0 0 0 0 0 ,

t T t T T TT

Heat and press cycle (c) Diffusion bonding of plasmasprayed monolayer tapes

Idooooooool Consolidated foil

Figure 5-36. Schematic representation of manufacturing methods for MMCs-I (Guruganus et al., 1988).

5.4 Main Aluminum Alloy Groups

5.4.11.3 Continuous Fiber Reinforced Al Composites

The most widely used production processes are (a), (b), and (c) as shown above. There are various kinds of continuous

Ejector Die/preform preheating

Molten-metal pouring #4

Thermocouple locations High-pressure infiltration Removal of composite (a) Squeeze casting

Direct extrusion

Evaluation

(b) Powder metallurgy method

Figure 5-37. Schematic representation of manufacturing methods for MMCs-II (Guruganus et al., 1988).

263

fibers suitable for incorporation with aluminum alloys such as 6061, 2014, 5456, and 7075, i.e., boron, silicon carbide, alumina, and graphite, as shown in Table 511. The relatively thick monofilaments of B and SiC made by chemical vapor deposition (CVD) are available to produce an acceptably reliable composite which has already been used. Graphite multifilament continuous fibers are made from two precursor materials, i.e., polyacrylonitrile (PAN) and pitch, and they offer a wide range of values of strength, modulus, and resultant cost which influence the overall cost of the composites, depending upon the degree of graphitization. There are also Japanese multifilament SiC fibers which are made by pyrolysis of organic compounds. Al composites have an excellent combination of properties, such as high specific strength and stiffness together with a lower coefficient of thermal expansion, high specific heat conductivity, and better wear resistance. The off-axis specific stiffness of a 6061 aluminum alloy composite is superior to those of poly(ether ether ketone) (PEEK) polymer matrix composites. For instance, the longitudinal tensile properties of MMCs developed in the R&D Project of Basic Technologies for Future Industries (in Japan) are summarized according to the relation between the specific strength versus the specific modulus and the temperature dependence of the tensile strength, as shown in Fig. 5-38 (Sakamoto, 1991). 5.4.11.4 Whisker or Short Fiber Reinforced Al Composites

The development of SiC whiskers or alumina-based short fibers, such as I d ' s SAFFIL fibers, and their incorporation into aluminum alloys by squeeze casting

264

5 Aluminum-Based Alloys

X10 2

O A D B T

10

(2 8 J2

to g 6

ft

1

C/Al SiC scs/Al

(A

SiC CVD/A1 SiC CVD/TJ

SiC (Whisker ) / A l

SCS-2/606KK/0.50, press) SCS-€/Ti-6AI-4V( V/ 0.49, press) I ™ 2/Al-4Ti( V, 0 49 nress) SCS-6/Ti-15M(h-5Zr-3Al( Vjr0.37, ' ^ ' ^ m Vl °A% W(S&) iw/, HIP) nir; | SCS-6/Ti-6M-2Sn-4Zr-6Mo (V/037, /~.M40J/1080(K/0.43, HIP) SCS-6/Ti-15V-3Cr-3Sn-3Al WJM40J/l080(V/0.43, press) IA1 A _ (Vj 0.35, press) -M40J/1080(V/0.44, roll) "1M4O/4O32(V/0.48, HIP) SiC(Nicalon)/1050 t M40/5056(Ky 0.47, press) (V, 0 40 rollL SCS-6/THA1-4V ( V / O i . p r e s s ^ ^ iV, 038, HIP) M40J/1080(V, 0.44, laser)

M40J/606KV/0.67,

O UHM/4032(V/0.56,HiP)

O

7075

6

8

10

12

14

16 X10 4

Specific modulus (MPa/r)

(B)

2 000

-2/Al-4Ti(V> 0.49, press) SCS-2/Al-8Cr-lFe( V> 0.48, press)

T Yi SCS-6/Ti-15Mo-5Zr-3Al 1\ \ (V> 0.37, HIP) SCS-6/Ti-15V-3Cr-3Sn-3Al (V> 0.35, press) M40J/1080( Vj 0.43, HIP) •M40J/1080( V) 0.43, press) HM40/4032(V> 0.40, HIP) N N ^iC(Nicalon) /1050( Vf 0.40, roll)

1000

Ti-6A1-4V'

7075-T6

'300

400

500

600

700

800

900

Temperature (K) Figure 5-38. (A) Specific strength versus specific modulus, and (B) temperature dependence of tensile strength of MMCs developed in R&D Project of Basic Technologies for Future Industries in Japan (Sakamoto, 1991).

5.4 Main Aluminum Alloy Groups

have been extensively exploited in their industrial applications, particularly in the automotive industry, where high-temperature properties, thermal fatigue, and wear resistance are important criteria, together with cost performance. The other important technology for incorporating whiskers or short fibers involves powder metallurgy techniques which blend the reinforcements with powder products made by RS or MA processes. The blends are then consolidated by various techniques into billets. A major advantage of the powder metallurgy approach is the ability to utilize the improved matrix properties resulting from the advanced processing mentioned earlier in the composites. The two technologies, squeeze-casting and power metallurgy, result in a different set of properties as well as cost requirements, and hence have to be taken into account depending on what is required. Applications such as aircraft structures require a damage-tolerant material which is more performance than cost sensitive. Other markets, such as automotive as well as industrial applications, tend to be much more cost sensitive and often require a different set of properties than the aircraft structure applications. It should be emphasized that both technologies have the added advantage of producing the final product with more or less isotropic properties, which is not possible in the case of continuous unidirectional reinforced composites. The billet materials produced by these two techniques can be fabricated using conventional working technologies such as extrusion, hot forging, and superplastic forming. The most successful example of an application is a diesel engine piston with an MMC-reinforced ring groove, as shown in Fig. 5-39. The squeeze-casting method is

265

Figure 5-39, Piston with MMC-reinforced ring groove developed by Toyota (Kaneko, 1991).

used to allow permeation of the aluminum alloy melt into a preformed reinforcement fiber such as SAFFIL-A12O3 (crystallized) or Kaowool A12O3 (amorphous) fibers. The present production volume is 400 000 units per month, which is quite a large volume and not found elsewhere in the world for MMC parts (Kaneko, 1991). 5.4.11.5 Particulate-Reinforced Al Composites Particulate-reinforced composites have been used industrially for many years because it is relatively easy to produce large quantities of these materials by conventional or slightly improved manufacturing methods, such as casting and various hotworking methods. The billets are made by introducing the particulate into molten or partially solidified metals, followed by the casting of these slurries into molds. Recently, Duralcan, a composite based on Al-Si alloys containing up to 20 vol.% in-

266

5 Aluminum-Based Alloys

expensive SiC grinding particulates, has been produced by a proprietary ingot metallurgy process by Dural Aluminum Composites Corporation. The Duralcan composite has a specific stiffness up to 45 % greater than that of steel, titanium, and other metals, a specific strength up to 50 % greater than that of aluminum, and its wear resistance is 10% greater. The potential advantage of this type of composite is the ability to produce near-net-shape components in a simple and cost-effective manner at a rapid manufacturing rate for large numbers of complex-shaped composite components (Duralcan USA, 1991). 5.4.11.6 Recent Applications of Al Metal Matrix Composites (MMCs) Since a significant improvement in the performance of MMCs has been widely recognized, MMCs have already been used in various industries such as automotive and leisure goods, as well as aerospace, and others (Murakami, 1991). Automotive applications (Rohatgi, 1991) tend to be much more cost-sensitive than aerospace applications, and hence it appears that discontinuous fiber composites incorporating particulates or short fibers using less expensive stir-casting or squeeze-casting techniques are most promising in this sector, although considerable attention has been devoted to powder metallurgy routes. In order to produce lighter, quieter, and more fuel-efficient automobile engines, it may be very effective to produce aluminum composites as standard components such as pistons, connecting rods, gudgeon pins, cylinder liners, and valve train parts. The unique properties obtainable with MMCs include high strength and stiffness together with better wear resistance, lower coefficient of thermal expansion, and high specific heat con-

ductivity. The most successful application of MMCs as automotive components is at present the selective reinforcement (using short fiber preforms) of the crown and ring groove in diesel engine pistons, as described earlier. The potential aerospace market for MMCs (Charles, 1990; Frazier et al. 1989; Wadsworth and Froes, 1989) is considered capable of becoming extremely large, and a recent survey predicts that the U.S.A. market will expand to over $ 200 M by the year 2000. The conventional aero-materials have almost reached their limits, and hence a significant proportion of future aero-engines are likely to be fabricated from MMCs. MMCs will be used in every sector of the aerospace industry such as aero-engines, airframes, spacecraft, and guided weaponry. It is well known that the space shuttle bay structures are made from boron monofilament reinforced aluminum tubes, where a 44% weight saving over aluminum alloys was achieved with an added advantage of lower thermal conductivity, resulting in reduced thermal insulation requirements. At present, a major interest in spacecraft has been aroused in manufacturing the truss structures from graphite fiber reinforced aluminum which can be made with a virtually zero thermal expansion coefficient. MMCs have already been used in various leisure and sporting sectors, similar to fiber-reinforced plastics for fishing rods, tennis rackets, golf club heads, etc. 5.4.12 Metallic Glasses and Nanocomposites Based on Aluminum Metallic glasses have been studied since 1959, when they were first prepared by RS. The most important glasses are those based on iron; some of these are today used on a substantial scale as transformer

5.5 Major Application Fields of Aluminum

laminations. Others are used for their corrosion resistance, electrocatalysis, and other properties. (See Vol. 9, Chap. 9, especially Sec. 9.6, for a treatment of metallic glasses and their uses.) In spite of the fact that metallic glasses can be very strong, there have been no serious attempts to exploit this characteristic until recently. This situation may now change because of the recent development of glasses based on aluminum, some containing more than 90 at.% Al. This family of glasses was discovered independently in Japan, France, and the U.S.A., as outlined by Cahn (1989). A vigorous development program in Japan since 1981 has established a number of glasses based on Al-ETM-LTM and especially Al-TM-Ln combinations, where TM are transition metals (ETM = early TM and LTM = late TM) and Ln are lanthanide metals. Examples are Al 70 Ni 20 Zr 10 , Al 90 Fe 5 Ce 5 , and Al 89 Ni 7 Y 4 , but there are numerous other good ternary and quaternary compositions incorporating Ni, Co, Fe, Y, La, Ce, etc. (Inoue and Masumoto, 1990; He et al., 1988). Such glasses often have yield stresses over 1000 MPa, though their elastic moduli are rather poor. They can be made as ribbons or wires (the latter by quenching into rotating water annuli). More recently it has been found that with appropriate compositions the material is partly nanocrystalline as-quenched, and such material is stronger, as well as stiffer, than pure glass. Thus Kim etal. (1991) have studied a series of Al88Y2Ni10_JCM:x. compositions, where M is Mn, Fe, or Co, and found that the nanocomposites in a glassy base thus prepared were ductile up to 2 at.% Mn, 5 at.% Fe5 or 7 at.% Co. Some of these nanocomposites had extraordinary yield stresses of up to ^ 1550 MPa with a 25% volume fraction of nanocrystals. Such good properties depend upon the presence of minute crystallites of a few nanometers

267

in diameter; crystallization of a pure glass gives coarse crystallites and the resultant material is brittle. He and his coworkers in America have found that in some ternary pure glasses, nanocrystals can be generated locally by plastic deformation of the ductile glass, and this may be another way of making high-strength materials. Applications of these intriguing materials can be expected soon.

5.5 Major Application Fields of Aluminum Table 5-12 illustrates the major application fields of aluminum materials, which have originated from their excellent intrinsic characteristics together with the great advantages of manufacturing and fabrication processes that are easily applicable for manufacturing the high-performance products described in the preceding sections. In the following sections, the major applications will be described briefly (see Brindenbauch, 1989; Hawkins, 1986; Nishi, 1992). 5.5.1 Packaging and Beverage Cans Aluminum foil and sheet consumption in packaging and beverage cans continues to expand. The excellent fabricating characteristics of aluminum allow the production of various kinds of containers by impacting, drawing, adhesive bonding, and spiral winding. Commercially pure aluminum is employed considerably in the form of foil or sheet, although nonheattreatable alloys of Al-Fe, Al-Mn, and A l Mg are also used for higher strength materials through varying degrees of work hardening. The three categories for the use of aluminum in packaging are described briefly in the following.

268

5 Aluminum-Based Alloys

Table 5-12. Major application fields of aluminum materials. Excellent intrinsic characteristics

Advantages in manufacturing and fabrication processes

Major application fields

Lightness High specific strength and modulus (excellent properties achievable from precipitation) High electric and thermal conductivity Good corrosion characteristics

low melting point easy melting and casting

beverage containers and packaging building materials

versatile deformation techniques applicable good machining

surface transportation industries; automotive and railroad cars aerospace industries

versatile joining techniques applicable advanced processing techniques applicable

electrical and electronic materials

Nonmagnetic, nontoxic, transparent to thermal neutrons

5.5.1.1 Flexible Packaging Flexible packaging is defined as involving the use of flexible materials such as foil or sheet to form a container. Aluminum foil is sheet that is less than 100 jam in thickness. Aluminum foil has various advantageous characteristics such as low permeability to water vapor and gases, tasteless, colorless, nontoxic, hygienic, effective light barrier, efficient reflector, good thermal conductivity, noncombustible, excellent dead-folding property, corrosion resistive to water, oil, grease, organic solvents, etc. Alloys used for foil in package applications include 1100, 1145, 3003, and 5052 with tempers such as O to HI9.

heat-exchangers other applications

has risen to almost 2000 cans per minute in a large-scale factory and the thickness of a beverage can body was reduced to 0.103 mm in 1990 from 0.132 mm in 1979, and its weight to 11.6 g from 15.5 g in the advanced plants, as shown in Fig. 5.40. Hence quality requirements of alloys used for cans, such as 3004 and 5182, are becoming more stringent because of the need for thinner walls for lower weight. Thus the alloy chemistry and heat-treatment, as well as the melting practice are being effectively improved. Moreover recycling of

1979

^-0.335 mm

1990 ^0.285 mm

-0.132 mm

-0.103 mm

0.432 mm

0.300 mm

5.5.1.2 Beverage Cans The beverage can industry is now the largest market in aluminum sheet. The aluminum beverage can has demonstrated continued growth by many cost-effective innovations, particularly in the U.S.A. and Japan. The recent productivity improvements in modern can plants are most impressive. For example (Kawashima, 1992), the production speed of 350 ml can bodies

Body 15,5 g End 5,2 g Sum 20,7 g

vtV

Body 11,60 g End 3,75 g Sum 15,35 g

Figure 5-40. Change in thickness and weight of body and end of 350 ml cans made from aluminum alloys from 1979 to 1990 (Kawashima, 1992).

5.5 Major Application Fields of Aluminum

used beverage cans (UBCs) is needed more urgently for the aluminum beverage industry, because UBCs represent an essential source of metal for the production of new can sheets. The amount of UBC recycling has increased significantly, e.g., to 62% in the U.S.A. and 42% in Japan in 1990. 5.5.1,3 Closures for Bottles and Collapsible Tubes

Aluminum closures are manufactured from 1100, 3003, and 3004 alloys most frequently using H14 and H34 tempers together with HI9 for 5052 just lately. Collapsible tubes are fabricated from slugs by impact extrusion of 99.5-99.7% commercially pure aluminum. The market is for toothpaste, medical ointments, cosmetics, shaving cream, adhesives, paint colorants, etc. 5.5.2 Building Applications

The building industry has for many years been one of the most important fields of application for aluminum. Extrusions for window sashes, sheets for panels, and castings for panels are standard nonstructural parts in modern large buildings as well as private houses. Architects and designers favor aluminum because of its advantages of light weight, ease of fabrication and erection, high precision, decorative appearance, and corrosion resistance. For extruded materials, 6000 series alloys are used mainly owing to their favorable extrusion properties, which originate from appropriate adjustments of the ratio of Mg to Si, homogenization treatment of the billets, and the cooling rate immediately after extrusion. The tensile properties of the extruded materials are then developed by simple age hardening. In the field of exterior materials, anodized materials enhance a favorable ex-

269

ternal appearance, which results from a smooth, uniform surface quality. This quality necessitates uniformity of the composition as well as the microstructure of the produced sheet. The aluminum materials for exterior applications are usually commercially pure aluminum of the 1050 and 1100 series, and the 5005 series, which afford a favorable quality after anodizing treatments and are compatible with 6063 extruded parts. Color-anodizing materials are very beneficial for exterior building uses, because no fading occurs due to sunlight. An anodized Al-Fe-Si alloy with additions of 1 % or more Fe in a sulfuric acid bath was reported to have a bluish, deep gray anodized color. 5.5.3 Surface Transportation Industries 5.5.3.1 Automotive Applications

The most prominent advantages of aluminum are its combination of light weight and adequate strength together with its resistance to corrosion, finishing characteristics, and easy maintenance. Hence aluminum is very favorable for designing a safer and cleaner car without adding too much weight for various construction parts and engines. Moreover, with growing environmental pressure to reduce noxious fume emissions and also to achieve maximum recycling of all car materials together with the increasing strategic importance of oil, aluminum seems to be significantly attractive, because its use leads to savings in car weight, which in turn leads to considerable fuel savings, and its low energy to remelt is significantly advantageous in terms of recycling costs. Aluminum is presently utilized for automobile engines in a larger tonnage than in bodies and chassis. Many wrought alloys have been developed for body sheets made from the slight-

270

5 Aluminum-Based Alloys

ly modified 2000, 5000, and 6000 series, which are called 30-30 alloys since they have 30 kgf/mm2 UTS (1 kgf/mm2 = 9.81 MPa) and 30% elongation. In this field, engine hoods, baggage doors, and fenders are the target parts. Since the modulus of aluminum is only one third that of steel, significant design changes are necessary to compensate for additional elastic deflection. Further development of materials with high strength, high formability, and high workability is urgently desired. Aluminum alloy cylinder blocks have become economically competitive with gray-iron castings since a die-cast design became available. The two prime factors that make aluminum attractive compared with cast iron are its lightness and its high thermal conductivity, which effects a significant reduction in the volume of coolant required in the system. Automotive radiators are mostly manufactured from aluminum in both Europe and the U.S.A., but not in Japan where the introduction of aluminum in this field was hindered by the existence of depreciated copper radiation production facilities together with cost and corrosion problems. However, aluminum heat exchangers with improved corrosion resistance are now being developed with considerable success. Further innovations not only of new materials, but also of brazing and other processing and manufacturing technologies, are expected in order to enlarge the share of aluminum in radiators. Aluminum wheels have become more common mainly because of their excellent function and appearance. Passenger-car wheels are manufactured as permanent mold castings from A356(Al-7% Si0.35% Mg)-T6; any alloy selected has to have sufficient ductility to prevent service failures due to careless driving.

Aluminum MMCs are currently receiving serious attention with respect to their higher specific strength and modulus, which may be beneficial in reciprocating components such as pistons, piston rings, valve gears, and connecting rods, as described in Sect. 5.5.4. The use of aluminum in commercial cars is based largely on the manufacturing costs, although engineering advantages and the economy ultimately determine whether aluminum will be used. In order to expand the aluminum market in automotive industries, the relative costs and availability have to move closer to those of rival materials if they are to make an impact on current design. For pistons in passenger cars, precipitation-hardened alloys with high contents of silicon (e.g., 12%) are used. The high content of silicon reduces the coefficient of thermal expansion and improves the wear resistance against the piston ring. 5.5.3.2 Railroad Cars

An Al-Zn-Mg alloy, 7N01 (7004), with excellent weldability and good natural age hardenability in the as-welded condition, was developed to produce aluminum bolsters for railroad cars in Japan. In addition, the development of a 6N01 alloy, which has good weldability, easy extrusivity, and corrosion resistance, made the production of floor structures fabricated from extruded hollow section members possible. 5.5.4 Aerospace Industries

Aluminum has been the main structural material for commercial and military aircraft for almost 70 years, and conventional high-strength wrought aluminum alloys still play an important role in aircraft and aerospace industries. However, high-performance aerospace is now creating a de-

5.5 Major Application Fields of Aluminum

mand for new aluminum materials, not only for airframe and engine applications, but also for missile and space systems. Recently, great advances in aluminum materials have been achieved using a variety of processing technologies, including ingot metallurgy, powder metallurgy, rapid solidification, mechanical alloying, and composite technology. Present applications of aluminum include airframes, landing gears and wheels, reciprocating and turbine engine components, propellers, systems elements, and interior trim. The advantages of aluminium for these applications are its lightness, high specific strength, and weatherability. For engine applications, its thermal conductivity is also a crucial property The conventional high-strength wrought alloys for airframes were: duralumin, an Al-Cu-Mg alloy which originated in Germany and was developed in the U.S.A. as alloy 2017-T4, which has 27 kgf/mm2 yield strength (YS) and is utilized primarily as sheet and plate. Then a higher strength alloy 2024-T3 with 35 kgf/mm2 YS was developed initially as alclad sheet. Alloy 2014-T6 was used for forging because of its significant artificial age-hardening ability. Lastly alloy 7075-T6, and A l - Z n - M g - C u alloy with 49 kgf/mm2 YS was introduced in 1943. Since then, most aircraft structures have been specified using this type of alloys. Many thousands of light planes are equipped with four-cylinder and six-cylinder air-cooled engines of an in-line reciprocating type employing cast aluminum crankcases of 355-type alloys (Al-5% Si1.25% Cu-0.5% Mg), and cylinder heads typically of 142-T7. The future super- and hypersonic aircraft, missiles, and space systems, are required to be lighter, with increased temperature capability, than the present generation of aluminum materials.

271

The development of advanced aluminum materials is briefly described in the following. An Al-Li alloy, 2020-T6, was introduced in 1957 for airframe applications, but soon afterwards its development was postponed because of its low fracture toughness. Recently, with the removal of this hinderance to development, the research has become very active again. The promising alloys contain up to 2.5wt.% Li, l-3wt.% Cu, 0-1.5 wt.% Mg, and 0.2 wt.% Zr. Al-Li alloys are described in Sect. 5.4.9. In contrast to the above I/M approaches, powder metallurgy/rapid solidification or mechanical alloying (P/M/RS or MA) aluminum-based systems have been investigated for the following three classes of alloys: (i) High-strength, precipitationhardended alloys, for example, X7064, 7090, 7091, and CW67; (ii) Low-density aluminum alloys of two groups: one based on extending the Al-Li alloy composition beyond the I/M limit for Li of about 2.7 wt.%, and another including substantial additions of Be, e.g., more than 10% to Al-Li I/M alloys; and (iii) Dispersoidstrengthened, elevated temperature alloys including alloys containing solute levels approaching 10-15 wt.% transition elements, such as A l - 8 F e - 2 M o , A l - 8 F e 4Ce, A l - 5 F e - 3 N i - 6 C o , Al-12Fe-2V, Al-4 Cr-3 Zr, etc. 5.5.5 Electrical and Electronic Materials For electrical conductor uses electrical conductor grade (EC) metal containing 99.6% Al with a conductivity of 62% IACS on a volume basis is commonly used with a UTS of 8.4-20.3 kgf/mm2 depending on the temper. For powder transmission and distribution lines, the main requirements are high electrical conductivi-

272

5 Aluminum-Based Alloys

ty, high mechanical strength and fatigue resistance, and good connections and joints. Hence both the heat-treatable alloy 6201-T81 (0.7% Si, 0.75% Mg) and the nonheat-treatable alloy 5005-H19 (0.8% Mg) have been employed as standard aluminum conductors. A great advantage of aluminum conductors is that at very low temperatures the electrical conductivity of high-purity aluminum is better and less impaired by high magnetic fields than that of copper. Hence high-purity aluminum is very suitable for cryogenic conductors for magnet coils of elementary particle accelerators, and as a sheathing material for superconducting Nb-Ti alloy wires. Aluminum is preferred for magnetic memory disk substrates as it is a reliable nonmagnetic material with excellent durability. The main alloy used is the 5086 alloy (4.0% Mg-0.45% Mn-0.15% Cr), which is manufactured with strict control of the quantity and size of coarse particles of Fe-Mn or Mg-Si compounds using melt treatment and filtration to eliminate inclusions. Other aluminum alloys are used, on account of their nonmagnetic characteristics, to make the tape feeding cylinder; one of the most important parts of a video recorder. The 'most important requirements are low abrasiveness and low friction during sliding contact with the magnetic tape, as well as good machinability and a fine surface finish. Conventional video recorder cylinders are manufactured from Al-Cu castings and Al-Si die castings, but they suffer from several disadvantages. Hence some new alloys such as cold-forged 2028 series alloys and the RSP/M alloy of Al-20% Si-2% Cu-1 % Mg are being introduced for improved abrasion resistance. The application of aluminum foils as electrodes for electrolytic capacitors is im-

portant in second place to packaging foils. Anode foils use high-purity aluminum of 99.95-99.995%, while cathode foils are manufactured from 1070, 1085, or 3003 alloys. To increase the capacitance, both the anode and cathode foil have to be subjected to surface roughening by electrochemical etching treatments. The dielectric in aluminum capacitors consists of oxide films formed on the etched surface of foils by anodic oxidation, the thickness of which varies in proportion to the chemical conversion voltage. Thus the surface area of oxide films formed by low-voltage chemical conversion can be increased by forming fine and complex pits, while foils for highvoltage chemical conversion necessitate relatively coarse pit formation. The surface roughening of anode foils for high-voltage chemical conversion is usually performed by direct current etching treatments. The efficency is affected by several factors, such as the purity of the aluminum base, additions of trace elements, recrystallized microstructures with a high volume fraction of crystal grains with cubic orientations, and an optimum oxide film thickness, which cause tunnelshaped etch pits which increase the surface area of oxide films. On the other hand, the surface of anode foils for low-voltage chemical conversion is roughened by alternating current etching. The area of the roughened surface is controlled by adjusting such conditions as the amount of additional elements, such as Fe, Si, Cu, and Mg, as well as reduction of cold-rolling and heat-treatments.

5.5.6 Other Applications The main applications are described in the preceding sections. In addition, aluminum is one of the preferred structural

5.6 Concluding Remarks

materials for radiation-resistant nuclear applications. Aluminum alloys were employed as ultrahigh-vacuum vessel materials for the gigantic particle accelerator "TRISTAN" due to their far more rapid attenuation of induced radioactivity, together with very little gas emission compared to stainless steels. Aluminum alloys are favorably adopted for nuclear fusion experimental apparatus in terms of weaker induced radioactivity and a much faster attenuation rate of generated radioactive nuclides. Several alloys such as Al-Mg-Si, Al-Mg-V, Al-Mg-Li, and SAP (sintered aluminum powder) are now being developed in relation to problems such as low residual induced radioactivity, heat and corrosion resistance, and irradiation damage.

5.6 Concluding Remarks 5.6.1 Future Applications and Possible Barriers

The future of aluminum and its alloys involves the large field of structural as well as functional materials. The development of advanced materials has always been promoted by two, often opposite requirements; one is to pursue improved performance which gives rise to higher cost, and the other is to obtain a reduction in the cost of the materials or components in order to ensure overall competitiveness among the candidate materials for industrial applications. To analyze the future applications we have to have a clearer conception of the manner of living of people in the next century. Each country has its own particular circumstances in the demand for aluminum materials. For instance, the U.S.A. has a large aerospace-related industry which places more emphasis on per-

273

formance than on cost sensitivity, while Japan maintains civilian-related industries such as automobiles, which tend to be much more cost sensitive. In order to breakdown the possible barriers, it may prove to be more important to meet the users' social and economic needs. The main pointers to the future seem to be as follows: (1) A key element is the continued development of the fundamental understanding of the microstructure-property relationship and the process-microstructure relationship in order to raise their technological level, to develop alloys with higher added value, and to produce them more efficiently. The development of new alloys has always involved a struggle between the two often conflicting requirements: improved performance and a reduction in cost to improve the overall competitiveness of the product. (2) As for existing alloys, more improvements in the structural as well as the functional qualities may be demanded, together with reasonable cost savings. To satisfy these requirements, it will be necessary to develop high-quality but low-cost alloys, which can only be produced using mass production technologies such as novel forming processes, new bondings, and improved surface treatments. For example, aluminum alloys for beverage cans are now the main sheet application field, which will be expanding, particularly in Japan, in the future. The main problem is cost reduction, which can be realized by further reduced weight and UBC recycling. Weight reduction can be achieved by improving the alloy properties, as well as by better sheet manufacturing and canforming technologies. (3) The demand for advanced airframes and propulsion systems requires the development of structural materials with prop-

274

5 Aluminum-Based Alloys

erties largely superior to those which are now available. New processing technologies and composite materials are being developed in the aluminum alloy field. Alloys produced by rapid solidification and mechanical alloying technologies lead to unique properties that are not otherwise attainable. The precise surface treatment technologies are very effective at improving surface-related characteristics such as surface strength, corrosion resistance, and lubrication. Moreover, superplastic forming, isothermal forging, and CIP and HIP powder metallurgy technologies offer net or near-net shape components which result in significant cost savings. Technology advances are certainly playing a large role in cost saving for advanced materials. (4) Competition and substitution among the different materials exerts a serious influence on the application fields for each material. There are some examples, for instance, glass fiber reinforced plastics replace certain aluminum die castings. Future material substitution has to be evaluated by the overall cost-effectiveness, which involves the cost of raw material processing to finished end product, the ease and safety of subsequent maintenance, performance in service, and moreover the recycling cost with environmental considerations.

conservation of natural resources, as well as energy. Furthermore, in the future rapid scrap generation will take place because products produced in an age of rapid economic growth will require fairly rapid replacement. At present, the social problem of waste disposal, primarily in urban regions, is raising many problems such as the disposal of scrap cars, used cans, residual ash, etc. Therefore considerable effort has to be made to establish advanced technologies matched to present-day needs for the recovery, separation, and sorting of scrap, labor-saving melting, metal yield improvement, iron removal by physical or metallurgical methods, refining of magnesium and nonmetallic inclusions, and environmental preservation.

5.7 Acknowledgements The author sincerely thanks Professor R. W. Cahn and Dr. K. H. Matucha for their great assistance and various suggestions during the preparation of the manuscript. Section 5.4.12, Metallic Glasses and Nanocomposites Based on Aluminum, was prepared by Professor Cahn, and Sec. 5.3.5.2, Mechanical Alloying Process, was rewritten by Dr. Matucha. These contributions are also gratefully acknowledged.

5.6.2 Aluminum Recycling

Global environmental protection is the critical issue. Highly important requirements are the restriction of CO 2 emissions, energy savings, the development of new energy, and the treatment and disposal technologies for industrial wastes. Development of the technologies for the recycling of useful metals, particularly aluminum, is urgently required to achieve global environmental preservation and the

5.8 References Adachi, M. (1984), J. Jpn. Inst. Light Met. 34, 430436. Arzt, E., Schultz, J. (Eds.) (1989), New Materials by Mechanical Alloying Techniques. Oberursel: DGM. Bagaryatskii, I. A. (1948), Zh. Tech. Fiz. 18, 827. Balmuth, E. S., Schmidt, R. (1980), in: AluminumLithum Alloys: Sanders, T. H., Jr., Starke, E. A., Jr. (Eds.). London: Met. Soc, p. 69. Benjamin, J. S. (1970), Met. Trans. 1, 294.

5.8 References

Blackburn, L. B., Starke, E. A., Jr. (1987), in: LightWeight Alloys for Aerospace Applications: Lee, E. W., Kim, J. J. (Eds.). Warrendale, PA: TMS, p. 110. Brindenbauch, P. R. (1989), Aluminium 65, 111. Cahn, R. W (1989), Nature 341, 183. Charles, D. (1990), Met. Mater. 42, 78-82. Chen, J. M., Sun, T. S., Viswanadham, R. K., Green, J. A. S. (1980), Metall. Trans. 11A, 85. Clyne, T. W., Withers, P. J. (1993), An Introduction to Metal Matrix Composites. Cambridge: Cambridge University Press. Danh, N. C , Rajan, K., Wallace, W. (1983). Metall Trans. 14 A, 1843. DiRusso, E., Conserva, M., Burrati, M., Gatto, F. (1974), Mater. Set Eng. 14, 23-26. Doig, P., Edington, J. W. (1975), Metall. Trans. 6A, 943. Duralcan USA (1991), Data brochures on Duralcan composites, Division of Alcan Aluminum Corp. Embury, J. D., Nicholson, R. B. (1965), Ada Met. 13, 403. Frazier, W. E., Lee, E. W, Donnellan, M. E., Thompson, J. J. (1989), J. Met. 41 (May), 22-26. Froes, F. H., Pickens, J. R. (1984), /. Met. 36 (Jan), 14. Froes, F. H., Kim, Y.-W, Herman, F. (1987), J. Met. 39 (August), 20-33. Gayle, E. W, Vander Sande, J. B. (1984), Scr. Met. 18, 473. Gest, R. I , Troiano, A. R. (1974), Corrosion-NACE 30, 214. Grimes, R. (1990), in: Supplementary Volume 2 of the Encyclopedia of Materials Science and Engineering: Cahn, R.W. (Ed.). Oxford: Pergamon, pp. 667679. Grimes, R., Stowell, M. X, Walts, B. M. (1976), Met. Technol. 3, 154. Gruhl, W. (1963), Z. Metallkd. 54, 86. Guinier, A. (1938), Nature (London) 142, 569. Guruganus, T. B., Gilliland, R. G., Hunt, H. (1988), in: Proc. Int. Symp. on Basic Technologies for Future Industry. Tokyo: Jpn. Ind. Tech. Assoc, pp. 99-142. Hardy, H. K., Heal, T. J. (1954), Prog. Met. Phys. 5, 143-278. Hawkins, J. E. (1986), in: Aluminium Technology '86, Proc. Int. Conf: pp. 38-42. He, Y., Poon, S. J., Shiflet, G. J. (1988), Science 241, 1640. Hellawell, A. (1990), in: Supplementary Volume 2 of the Encyclopedia of Materials Science and Engineering: Cahn, R. W (Ed.). Oxford: Pergamon, pp. 691-697. Higashi, K., Nagai S., Maeda, M., Ohnishi, T. (1986), /. Jpn. Inst. Light Met. 36, 361. Holve, K., Weber, W, Wincierz, P. (1968), Aluminium 44^ 467-475. Horiuchi, R., Kaneko, J. (1976), J. Jpn. Inst. Light Met. 26, 327.

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Humphreys, F. J. (1979), Met. ScL 13,136; (1980), in: Proc. 1st Riso Int. Symp. on Recrystallization: Hansen, N., Jones, A. R., Lefters, T. (Ed.). Ris0, Denmark: Ris0 Natl. Lab., p. 35. Inoue, A., Masumoto, T. (1990), in: Supplementary Volume 2 of the Encyclopedia of Materials Science and Engineering: Cahn, R.W. (Ed.) Oxford: Pergamon, pp. 660-667. Islam, M. U., Wallace, W. (1983), Met. Technol. 10, 386; 11, 320. Kaneko, R. S. (1980), Met. Prog., April, 41. Kaneko, T (1991), in: Proc. 2nd Int. Conf. on Advanced Materials and Technology, New Composites '91, Hyogo, Kobe, Japan, pp. 165-172. Kawashima, T. (1992), J. Jpn. Inst. Light Met. 40, 856. Kim, Y. H., Inoue, A., Masumoto, T. (1991), Mater. Trans. JIM 30, 599. Kinzey, D. (1978), Light Met. 2, 227. Kosuge, H., Takada, H. (1979), /. Jpn. Inst. Light Met. 29, 64. Lee, E. W, Frazier, W. E. (1988), Scr. Met. 22, 53. Lorimer, G.W., Nicholson, R.B. (1966), Acta Met. 14, 1008-1013. Martin, I W (1968), Precipitation Hardening. Oxford: Pergamon. Matsuki, K., Uetani, Y, Yamada, M., Murakami, Y. (1976), Met. Sci. 10, 235. Metals Handbook (1967), Vol. 3. Metals Park, OH: ASM Int., p. 440. Metals Handbook (1973), Vol. 8. Materials Park, OH: ASM Int., p. 262. Murakami, Y (1981), J. Jpn. Inst. Light Met. 31, 748. Murakami, Y (Ed.) (1985), Fundamentals and Industrial Technologies of Aluminum Materials (in Japanese). Tokyo: Japan Light Metals Association. Murakami, Y (1990), Trans. JIM 31, 669-678. Murakami, Y (1991 a), in: Proc. RASELM '91: Hirano, K., Oikawa, H., Ikeda, K. (Ed.). Tokyo: Japan Inst. Light Metals, pp. 3-10. Murakami, Y (1991b), Proc. 2nd Int. Conf on Advanced Materials and Technology, New Composites '91, Hygo, Kobe, Japan, pp. 113-122. Nes, E. (1980), in: Proc. 1st Riso Int. Symp. on Metallurgy and Materials Science: Hansen, N., Jones, A. R., Lefters, T. (Ed.). Riso, Denmark: Riso Natl. Lab., p. 36. Nes, E. (1986), in: Microstructural Control in Aluminum Alloys: Chia, E. H., McQueen, H. J. (Ed.). Warrendale, PA: Metallurg. Soc, pp. 95-108. Nishi, H. (1992), Aluminium 68, 681-686. Nishio, T, Nasu, S., Murakami, Y (1970), J. Jpn. Inst. Met. 34, 1173-1178. Orowan, E. (1948), in: Symp. on Internal Stress. London: Inst. Metals, p. 451. Pashley, D. W, Rhodes, D. W, Sendorek, A. (1966), /. Inst. Met. 94, 41-49. Phillips, H. W L. (1946), /. Inst. Met. 72, 151.

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Phillips, H. W. L. (1959), Annotated Equilibrium Diagrams of Some Aluminum Alloys Systems, Monograph No. 25, London: Inst. Metals, p. 57. Pickens, J. R. (1990), in: Supplementary Volume 2 of the Encyclopedia of Materials Science and Engineering: Cahn, R. W. (Ed.). Oxford: Pergamon, 679-683. Pickens, J. R., Christodoulou, L. (1987), Metall. Trans. 18A, 135. Pink, E., Grinberg, A. (1984), Aluminium 60, S. 700. Polmear, I. J. (1989), Light Alloys, 2nd ed. Sevenoaks, U.K.: Edward Arnold, p. 45. Preston, G. D. (1938), Proc. Roy. Soc. A, Lond. 167A, 526. Rabiukin, A., Liebermann, H. H. (1993), in: Rapidly Solidified Alloys: Liebermann, H. H. (Ed.). New York: Marcel Dekker, 691-735. Rajan, K., Wallace, W, Beddoes, J. C. (1982), /. Mater. Sci. 17, 2817. Rohatgi, P. (1991), J. Met. 43, 10-15. Rosenstein, A. H. (1982), J. Met. 34, 14. Sakamoto, A. (1991), in: Proc. 2nd Int. Conf on Advanced Materials and Technology, New Composites '91, Hyogo, Japan, pp. 165-172. Savage, S. J., Froes, F. H. (1984), J. Met. 36, 20-33. Scamans, G. M. (1980), Metall. Trans. 11 A, 846. Scamans, G. M., Alani, R., Swann, P. R. (1976), Corros. Sci. 16, 443. Schmalzried, H., Gerold, V. (1958), Z. Metallkd. 49, 291. Silcock, J. M., Heal, T. I, Hardy, H. K. (1953-1954), /. Inst. Met. 82, 239. Slevolden, S. (1974), Met. Mater. 6, 94. Staniek, G., Wirth, G., Bunk, W. (1980), Aluminium 56, 699-702. Starke, E. A., Jr., Lingering, G. (1979), in: Proc. ASM Materials Science Seminar, p. 208. Structures and Properties of Aluminum (in Japanese) (1991). Tokyo: Japan Inst. Light Metals. Suzuki, H., Kanno, M., Hayashi, N. (1981, 1982), J. Jpn. Inst. Light Met. 31, 122; 32, 88, 577. Tamura, M., Mori, M., Nakamura, M. (1970), /. Jpn. Inst. Met. 34, 919. Theler, J. X, Furrer, P. (1974), Aluminium 50, 467. Thomas, A. T. (1966), Acta Met. 14, 1363-1374. Viswanadham, R. K., Sun, T. S., Green, J. A. S. (1980), Metall. Trans. 11A, 85.

Wadsworth, I, Froes, F. H. (1989), J. Met. 41 (May), 12-19. Waldman, J., Sulinski, H., Markus, H. (1974), Metall. Trans. 5, 573-584. Warlimont, H. (1977), Aluminium 53, 161. Watanabe, R. (1988), Bull. Jpn. Inst. Met. 27, 799801. Wert, J. A., Paton, N. E., Hamilton, C. H., Mahoney, M.W. (1981), Metall. Trans. 12A, 1267. Yoshida, H., Hashimoto, H., Yokota, Y (1982), Zairyo-kakaku 18, 361.

General Reading Altenpohl, D. (1965), Aluminium and Aluminium Legierungen. Berlin: Springer. Altenpohl, D. (1982), Aluminum Viewedfrom Within. Diisseldorf: Aluminium Verlag. Aluminium-Taschenbuch (1974), 13. Auflage: Nielsen, H., Hufnagel, W, Ganoulis, G. (Eds.). Diisseldorf: Aluminium-Zentrale. Das, S. K. (1993), "Rapidly Solidified P/M Aluminum and Magnesium Alloys - Recent Developments" in: Rev. Particulate Mater. 1, 1-40. Hatch, J. E. (Ed.) (1984), Aluminum - Properties and Physical Metallurgy. Metals Park, OH: American Society for Metals. Inoue, A., Masumoto, T. (1992), in: Conf. Proc. of ICAA 3, Vol. Ill, Trondheim, Norway, 22-26 June: Arnberg, L., Lohne, O., Nes, E., Ryum, N. (Eds.). Trondheim: Norwegian Institute of Technology, pp. 45-88. Loffler, H. (1994), Structure and Structure Development in Al-Zn Alloys. Berlin: Akademie Verlag. Metals Handbook (1975), Vol. 1, 8th ed. Metals Park, OH: American Society for Metals, pp. 865-958. Mondolfo, L. F. (1971), Metall. Rev. 16, 95-124. Mondolfo, L. F. (1976), Aluminum Alloys - Structure and Properties. London: Butterworths. Polmear, I. J. (1989), in: Light Alloys - Metallurgy of the Light Metals: London: Edward Arnold, pp. 15126. Van Horn, K. (Ed.) (1967), Aluminum, Vols. 1-3. Materials Park, OH: American Society for Metals.

6 Copper-Based Alloys Wolfram Heller Department of Electrical Engineering, Fachhochschule Miinchen, Munich, Germany

List of Symbols and Abbreviations 6.1 General Review of Copper and Copper Alloys 6.1.1 Alloying of Copper 6.1.2 Wrought Products 6.1.3 Cold-Worked Conditions 6.1.4 Most Important Properties 6.1.5 Fabrication and Deformation 6.1.6 Annealing and Hot Working 6.2 Technology and Requirements of Copper 6.2.1 Technology of Copper Production 6.2.1.1 Winning and Production of Copper 6.2.1.2 Melting and Refining 6.2.1.3 Pure Copper Products 6.2.2 Classification of Unalloyed Copper Materials 6.2.2.1 Tough-Pitch Oxygen-Containing Copper 6.2.2.2 Deoxidized Copper 6.2.2.3 Oxygen-Free Copper with High Conductivity 6.2.3 Copper Standards 6.2.3.1 Deutsches Institut fur Normung (DIN) Standards 6.2.3.2 American Society for Testing and Materials (ASTM) Standards 6.2.4 Copper Consumption 6.2.5 Requirements with Regard to Copper and Its Alloys 6.3 Properties and Behavior of Copper 6.3.1 Properties of Unalloyed Copper 6.3.2 Low-Alloyed Copper 6.3.2.1 Nonhardening Copper Alloys 6.3.2.2 Hardening Alloys 6.3.3 Conductive Bronze 6.4 Properties and Behavior of Copper Alloys 6.4.1 Copper-Zinc Alloys (Brasses) 6.4.1.1 Deformable Brass 6.4.1.2 Cast Brass 6.4.1.3 Special Brasses 6.4.2 Copper-Nickel-Zinc Alloys (Nickel-Silvers) 6.4.2.1 Deformable Alloys Materials Science and Technology Copyright © WILEY-VCH Verlag GmbH & Co KGaA. Allrightsreserved.

280 282 282 283 283 284 285 285 285 286 286 286 288 288 288 290 290 291 291 292 292 294 295 297 299 301 301 302 302 303 307 307 309 309 310

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6.4.2.2 6.4.3 6.4.3.1 6.4.3.2 6.4.3.3 6.4.4 6.4.5 6.4.5.1 6.4.5.2 6.4.5.3 6.4.5.4 6.4.5.5 6.4.6 6.4.6.1 6.4.6.2 6.4.6.3 6.4.6.4 6.5 6.5.1 6.5.1.1 6.5.1.2 6.5.1.3 6.5.1.4 6.5.1.5 6.5.2 6.5.2.1 6.5.2.2 6.5.2.3 6.5.2.4 6.5.2.5 6.5.3 6.5.3.1 6.5.3.2 6.5.3.3 6.5.4 6.5.5 6.5.6 6.5.7 6.5.8 6.5.8.1 6.5.8.2 6.5.8.3 6.6 6.6.1 6.6.2

6 Copper-Based Alloys

Cast Alloys Copper-Tin Alloys (Bronzes) Deformable Alloys Cast Alloys Red Castings Copper-Aluminum Alloys Copper-Nickel Alloys Special Electrical Applications Corrosion-Resistant Alloys Deformable Alloys Resistor Metals . Cast Alloys Copper Alloys with Silicon, Beryllium, Manganese, or Lead Silicon Bronze Beryllium Bronze Manganese Bronze Lead Bronze Technology and Products of Copper-Based Alloys Applications in Electrical Engineering and Electronics Electrical Conductors Construction Materials Contact Materials Resistor Metals Other Electrical Applications Copper Semiproducts Plates and Strips Tubes Rods and Profiles Wires Hot Forgings Copper Castings Conducting Materials Bearing Materials Construction Materials Sintered Copper Fabrication of Products by Non-Cutting Techniques Heat-Treatment and Annealing Metal Cutting Recycling of Copper-Based Materials Classification Electrical Scrap Automotive Scrap Copper Materials for Special Applications Corrosion-Resistant Materials Copper Coatings

310 311 311 312 313 313 314 315 316 317 317 317 317 317 317 318 318 319 319 320 321 321 322 322 323 323 323 324 324 325 325 326 327 327 329 330 331 332 332 333 333 336 336 336 337

Contents

6.6.3 6.6.4 6.6.5 6.6.5.1 6.6.5.2 6.6.5.3 6.6.6 6.6.7 6.6.8 6.6.9 6.6.9.1 6.6.9.2 6.6.9.3 6.7 6.8

Construction Materials Mechanical Composite Materials with Copper Copper Alloy Bearing Materials Solid Sliding Bearings Multilayer Metallic Materials Sintered Bearings Decorative Materials Copper Metallization Shape Memory Alloys Other Applications of Copper Copper as an Alloying Element Chemical Compounds with Copper Copper in Medicine Directions for Future Applications of Copper References

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338 338 340 340 340 341 341 342 343 343 343 343 344 344 344

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6 Copper-Based Alloys

List of Symbols and Abbreviations A5, A 10 C

P

E G L

•KpO.2

t T T

z

fracture elongation specific heat capacity modulus of elasticity shear modulus electrical conductivity value surface roughness fatigue strength tensile strength yield strength time temperature melting temperature area reduction

Qe

coefficient of thermal expansion elongation electrical conductivity thermal conductivity density electrical resistivity

ASM ASTM b.c.c. BS CATH CDA CIDEC C.I.S. CVD DCB DHP DIN DKI DLP DRAM E-Cu EEPROM ETP f.c.c. F-Cu FRHC FRTP G

American Society for Metals, Materials Park, OH American Society for Testing and Materials, Philadelphia, PA body-centered cubic British Standard cathode copper Copper Development Association, London, New York Conseil International pour le Developpement du Cuivre, Geneva Commonwealth of Independent States chemical vapor deposition direct copper bonding deoxidized high phosphorus residue Deutsches Institut fur Normung, Berlin Deutsches Kupfer-Institut, Berlin deoxidized low phosphorus residue dynamic random access memory electrolytic copper electrically erasable and programmable read-only memory electrolytic tough-pitch face-centered cubic fire-refined copper fire-refined high-conductivity fire-refined tough-pitch sand-mold casting

a 5 x X Q

List of Symbols and Abbreviations

GC GD GK GZ HB HV IACS IC ISO KE L MRS OF OFE OFHC OFLP PHC PVD S SE SF SMD SRAM SW ULSI UNS VDI

continuous casting die casting chill casting centrifugal casting Brinell hardness Vickers hardness International Annealed Copper Standard integrated circuit International Organisation for Standardization cathode electrolytic liquid melt Materials Research Society, Pittsburgh oxygen-free oxygen-free electrolytic oxygen-free high-conductivity oxygen-free low phosphorus residue phosphorus-deoxidized high-conductivity physical vapor deposition welding oxygen-free electrolytic phosphorus-deoxidized (high residual phosporus) surface-mounted device static random access memory phosphorus-deoxidized (low residual phosphorus) ultralarge-scale integration Unified Numbering System Verein Deutscher Ingenieure, Dusseldorf

281

282

6 Copper-Based Alloys

6.1 General Review of Copper and Copper Alloys Copper and copper-based alloys are among the most commercially important metals because of their excellent properties, ease of manufacture, and numerous applications. They are normally used because of their excellent electrical and thermal (heat) conductivities, outstanding resistance to corrosion, and ease of fabrication. Generally, copper alloys are nonmagnetic and have medium values of strength, hardness, and fatigue resistance. Copper and bronze were used together with gold and silver in the historical development of mankind as the first metallic materials. With the development of electrical engineering, the demand for copper with good electrical conductivity experienced rapid growth. Copper and copper alloys are of considerable importance because of their special characteristics, i.e., high electrical conductivity (conducting material), high heat conductivity (heat exchangers, coolers, radiators, chemical construction), easy to form (wires, strips, pots, plates), good chemical resistance (chemical apparatus, overhead cables, rain gutters), nonferromagnetic nature (magnets, coils), and color (jewellery, works of art). Copper alloys are differentiated as bronzes, brass, red brass, and nickel-silver, and moreover, as formable alloys and as cast alloys. Copper and its alloys, especially for electrical applications, can be readily soldered and brazed, and many copper alloys can be welded by various gas, arc, and resistance methods. Copper alloys can be polished and deformed to almost any desired texture and surface. They can be plated, coated with metals or organic substances, and chemically colored to various finishes with different forms, thicknesses, and surfaces.

Copper alloys with specific colors are readily available for decorative parts. 6.1.1 Alloying of Copper

Unalloyed, pure copper is mostly used for cables and wires, electrical contacts, and a wide range of construction parts, which are required for their electrical conductivity. Copper is a soft material and has a facecentered cubic (f.c.c.) crystal structure up to its melting point of 1083 °C (Table 6-1). The purer it is, the higher the conductivity, but conversely, the softer and weaker it is. To increase the strength, to make the material cheaper, or to obtain other special properties, copper is alloyed and treated further. For example, by heat-treatment or by cold or hot deformation the properties of copper and copper alloys can be changed.

Table 6-1. Physical and technological properties of unalloyed copper (depending on treatment). Property Density Melting temperature Linear thermal expansion Electrical conductivity Thermal conductivity Modulus of elasticity Tensile strength Yield strength Fracture elongation Brinell hardness Specific heat capacity Electrical resistivity

Symbol Q

Value

8.90-8.96 g/cm3

Tm

1083

a

16.8-17.2

X

Unit

35-58

°C

10~6/K m/(Q mm2)

X

240-386

W/(m K)

E

100-130 200-360 100-250

kN/mm 2 N/mm 2 N/mm 2

^p0.2

^ 5

HB C

P

2 - 45 % 55-110 kp/mm2 0.38-0.45

J/(gK)

0.03-0.017 Q mm2/m

6.1 General Review of Copper and Copper Alloys

die free form

extruding pressing

cold • • • •

rolling bending stamping drawing

softening • forging • rolling • bending • drawing

283

hardening

• soft annealing •hardening •diffusion •stress-free annealing annealing •recrystallization

Manufacturing to the final form Figure 6-1. Outline of possible manufacturing processes of copper and copper-based alloys.

Low-alloyed copper is mainly applied for electrically conducting wires and cables, because of its high electrical conductivity. Copper alloys, certain brasses (CuZn), bronzes (Cu-Sn), and cupronickles (Cu-Ni) are used for automobile radiators, heat exchangers, home heating systems, panels for solar heating, and various other applications requiring rapid heat conduction. Because of their outstanding ability to resist corrosion, copper, brasses, some bronzes and cupronickels are used for pipes, valves, and fittings in water-carrying systems and pipelines with other waterbased fluids. An outline of the available treatments of copper and copper alloys is presented in Fig. 6-1. Such treatments can be combined or used sequentially. The semiproducts thus produced are then manufactured to their final form. 6.1.2 Wrought Products

Wrought products of all classes of copper alloys are used in many manufacturing processes because of their easy machining or cold deformation properties, depending

on their copper content. Wrought products are sometimes in competition with cast copper alloys, because of their good shape, corrosion resistance, and easy machining. Most wrought copper alloys are available in various cold-worked conditions. Copper alloys containing 1-6% Pb have a good machining grade and are used mostly in cutting fabrication. 6.1.3 Cold-Worked Conditions

Copper and most copper alloys can be cold worked easily when they contain a large proportion of copper, and are in the a solid-solution range, because of their face-centered cubic (f.c.c.) crystal structure. On the other hand, copper alloys with another solid-solution structure or with precipitation are good for machining processes because of their more "brittle" crystal structure. Very heavy reductions are possible, especially with continuous flow of the material. Rolling reductions exceeding 90% in one pass are used for rolling strips. The result is often a preferred crystal orientation or texturing. Textured metals have dif-

284

6 Copper-Based Alloys

ferent mechanical properties in different directions (anisotropy), which is undesirable for some applications. Typical applications of cold-worked materials include springs, fasteners, hardware tubes and plates, and small gears and cams. Certain types of (mostly) small parts are produced by hot forging, e.g., plumbing fittings and valves. 6.1.4 Most Important Properties Copper and its alloys are mostly chosen instead of other metals because of their good electrical and thermal conductivity, good resistance to corrosion, special color, and ease of fabrication by either cold working or machining (Fig. 6-2). Other criteria for their selection are strength, resistance to fatigue, and ability to form a good finish, but these are of lesser importance. Copper resists most corrosive conditions quite well, although its alloys are sometimes limited in certain environments and corrosion protection is required because of hydrogen embrittlement or stresscorrosion cracking. Annealing or stress relieving after forming frequently prevents stress-corrosion cracking.

Cu

Eu-Zn.36

Cu-Sn8-(P)

Copper and its alloys are good conductors of electricity and heat. Therefore they are used more often than any other metal. Alloying decreases the electrical conductivity and, to a lesser extent, the thermal conductivity too, but increases the strength and mechanical properties (Table 6-2). For this reason, copper and high-copper alloys (low-alloyed copper) are preferred over either pure copper (lower strength) or alloyed copper (lower conductivity). The small amount of reduction in the conductivity due to low alloying can be improved by annealing. For decorative purposes, or when a special color and finish are required together with other properties, copper and certain copper alloys are used, i.e., soft pink copper, red-brown bronze, yellow-gold brass, gray-white nickel-silver. High purity copper is very soft, especially when undeformed and with few impurities, because the copper matrix is not then hardened by foreign atoms. By alloying and cold working, the tensile and yield strengths and the hardness can be increased to higher values. For copper and many copper alloys, the tensile strength in the hardest cold-worked condition is approximately twice that of the annealed

Cu-Be2 Cu-Co2,5-Be0.5 Cu-Ni25-Zn 15

Electrical conductivity,^ Tensile strength ,(RJ Spring stability Heat resistance Basic material cost Figure 6-2. Comparison of some properties of typical copper alloys (Heller, 1976).

6.2 Technology and Requirements of Copper

285

Table 6-2. Comparison of electrical conductivity x, tensile strength Rm, and material price of copper alloys in relation to unalloyed copper (= l) a . Compared values

Cu

Cu-Zn36

Cu-Sn8-(P)

Cu-Be2

Cu-Co2.5Be0.5

Cu-Ni25Znl5

Relative electrical conductivity (x) Relative tensile strength (Rm) Relative price

1 1 1

0.33 1.5 0.6

0.14 2.5 0.8

0.33 4.5 2.4

0.67 2.5 1.8

0.08 1.6 1.3

x/price

1 1

0.55 2.5

0.18 3.1

0.14 1.9

0.37 1.4

0.06 1.2

^Jprice Heller (1976).

temper, and the yield strength may be as much as five to six times higher. The relation between hardness and strength is different for different copper alloys. The hardness and strength can usually only be correlated over a narrow range of conditions. 6.1.5 Fabrication and Deformation Generally, copper and its alloys are capable of being shaped to the required form and dimensions by any of the common fabricating processes. All copper forms are normally rolled, stamped, drawn, and cold deformed, and at elevated temperatures they are rolled, extruded, forged, and formed. Casting is used for all copper alloys, especially small pieces. Coarsegrained copper alloys are somewhat softer than the fine-grained metals and are therefore somewhat more easily cold worked and machined (ASM, 1979). 6.1.6 Annealing and Hot Working By heating or annealing, the work-hardened metals can be returned to a soft condition as a result of recovery, recrystallization, and grain growth (see Vol. 15, Chap. 9 of this Series). The grain size is smaller and more uniform when deformed copper alloys are recrystallized. Large

grain sizes are normally produced by a combination of limited deformation and a long annealing time. High amounts of prior cold working, fast heating to the annealing temperature, and short annealing times favor fine grain sizes. This results in higher strength and, to a small amount, in lower conductivity. While the variation of the mechanical properties obtained by cold working is high, the variation of the mechanical properties by annealing is rather small, but by no means negligible. Fine grain sizes are often required to enhance the condition of annealed or cold-formed end products, such as higher strength, fatigue resistance, resistance to stress-corrosion cracking, or polished surface quality.

6.2 Technology and Requirements of Copper The properties of copper and its alloys greatly depend on the amount of alloying elements and even on small additions (traces) of impurities, hence the method of producing the copper from its ores is very important and decisive for the final application. The products and properties of pure and alloyed copper are therefore nationally and internationally standardized.

286

6 Copper-Based Alloys

As such, manifold requirements are imposed on the copper material, depending on its method of production, its semiproducts, its content of alloying elements, and the technology and treatment of products during final manufacturing. 6.2.1 Technology of Copper Production 6.2.1.1 Winning and Production of Copper Copper ores are found in many parts of the earth. They are different in kind, composition, and copper content. The main deposits of copper minerals are located in Chile, the U.S.A., Canada, C.I.S. (Soviet Union), and Zambia. Two different types of ores are mined (oxidic and sulfidic ores) and processed by different methods, together with recycled material, to raw copper containing ~ 99 % Cu 1 . Oxidic ores are processed hydrometallurgically by older methods using sulfuric acid as a solvent, or using the modern technology of solvent extraction followed by a refining process, i.e., electro winning. The final product obtained by electrolysis is cathodic copper (SE-CU) (SE: oxygen-free electrolytic). Sulfidic ores are processed pyrometallurgically by roasting and blowing the melt to fire-refined copper with some impurities. This impure copper ( ~ 9 9 % Cu) can be used directly or further refined by electrolysis to produce electrolytic copper (ECu) containing >99.9% Cu. 6.2.1.2 Melting and Refining Copper is nowadays produced from four types of raw material: 1) sulfide concentrates (for about 90% of copper production), 1

In this chapter, all contents of copper and copper alloys are given in weight % (mass %).

2) oxide and silicate ores (with lower copper contents), 3) bituminous ores, and 4) scrap metal and other secondary materials. Concentrates of ~ 15-35% Cu are produced by flotation from ores containing 0.5-2% Cu, and mainly consist of chalcopyrite (CuFeS2). However, oxide, silicate, and certain bituminous ores cannot be concentrated by flotation with good recovery and are therefore processed directly. The result is that there are two different technological and logistic processes: sulfidic concentrates are processed exclusively by pyrometallurgical methods (smelting) with energy, while chiefly oxidic ores are processed hydrometallurgically (leaching). Because there is a lot of gangue, the processing of nonfloated ores must necessarily be located close to the mine and demands much energy to melt the ore and the gangue. Hence the transportation of oxidic ores is uneconomic. Pyrometallurgical Process For pyrometallurgical processes, the melting and refining processes used for mainly sulfidic raw materials are basically identical and are summarized in Fig. 6-3. Raw ore materials are broken, ground, and floated to make concentrates and then melted autogenously in a reverberatory furnace to an intermediate product, called copper matte, by using the heat generated by sulfur reduction (roasting). Afterwards, the liquid copper matte (copper-iron-sulfide) is blown with air in different processes (slag reaction and blowing) to make raw or blister copper, until most of the iron enters the slag and most of the sulfur is burnt out. Blister copper contains small amounts of iron, nickel, arsenic, sulfur,

6.2 Technology and Requirements of Copper

Copper intermediate product

Copper content (wt.-%)

ements from the mining location and the origin of the scrap. This copper is suitable for applications requiring low electrical conductivity. The pyrometallurgical development of (impure) copper concentrates at present depends essentially on optimization of the off-gas streams and energy consumption.

Production process Mining

Sulfidic ores

287

0.5-2 %

Grinding, Flotation Concentrate

Hydrometallurgical Process

Reduction of sulfur (roasted) Reverberatory furnace Copper matte

40-75 °/e Converter (slag blow)

Raw (blister) copper

98-99 % .

Fire-refined (anode) copper

. JJ Anode furnace (pyremttaUurgicat refminq}

.

98.5-99.5%

Electrotysis (electrolytic refining) Electrolytic (cathode) copper

The hydrometallurgical process is at present used for around 10 % of all copper and has been growing qualitatively since 1970. This is due to the introduction of solvent extraction technology followed by electrowinning, the pure copper end product of which is equal in quality to that refined pyrometallurgically in a furnace. The most used solvent for oxidic or sometimes mixed oxidic-sulfuric ores is sulfuric acid (Metallgesellschaft, 1993). This new, combined technology also gives an economic process of leaching residues, which still have a low copper content of less than 0.5%.

to 99.98 % Remelting or poling (tough pitch)

Semiproducts: pigs, blocks, wire bars, round and flat rolling bars

Figure 6-3. Technology of copper production. Melting and refining the raw material of mainly sulfidic ores (Armstrong and Smith, 1970).

and noble metals, and is often melted by pyrometallurgical refining in a reverberatory furnace, sometimes together with copper scrap (Fig. 6-36). Fire-refined (anode) copper is heavily influenced by the amount of additional el-

Copper Refining Copper refining is achieved either by pyrometallurgical refining in anode furnaces to make fire-refined copper (anode copper) or by electrolysis to make electrolytic copper. During converter blowing (Fig. 6-3), over-oxidation of raw (blister) copper occurs, so that the copper generally contains over 0.5 % oxygen, which leads to a severe reduction in the conductivity. Afterwards the oxygen content is reduced to less than 0.2%, either by poling of liquid blister copper, or remelting of solidified blister copper followed by poling. With the poling process, oxygen reduction is carried out in rotary furnaces with (originally) logs (birch wood) or with natural gas, naphtha

288

6 Copper-Based Alloys

or reforming gas through the formation of volatile organic compounds. By remelting, other materials such as scrap or the remains of anodes are then melted together with the blister copper, often in a reverberatory furnace. 6.2.1.3 Pure Copper Products Fire-Refined Copper Fire-refined copper (F-Cu) is used for applications in electrical engineering and further processing of the melt products to form plane surfaces. On solidification of poled or tough-pitch copper (tough with good forgeability), some gas will be freed, which compensates the volume shrinkage and results in flat casting surfaces. Electrolytic Copper In order to produce electrolytic copper of high conductivity, fire-refmed copper is further processed by electrolysis with a sulfuric acid based electrolyte in the cells. The anode contains casting plates of fire-refined copper ( ~ 9 9 % Cu), while the cathode consists of a thin starting sheet of pure copper which is negatively charged. Highpurity copper with an impurity content of only 10-20 ppm is deposited on the cathode under special process conditions, while base impurities such as nickel or arsenic remain in the electrolyte, and noble elements (Au, Ag, Pt) and insoluble compounds (Pb, Se) stay in the anode slime. Electrolytic copper with more than 99.9% Cu (OF-Cu) (OF: oxygen-free) can also be remelted as tough pitch and cast for various applications. Cathode copper for the highest requirements contains about 99.99% Cu(Cu-K, KE-Cu, Cu-OF).

6.2.2 Classification of Unalloyed Copper Materials The melting and refining produces two types of pure copper depending on the refining process: fire-refined copper or electrolytically refined copper. For applications, distinction of the characteristics and main properties is more important, depending on the treatment of the melted copper before casting or forming semiproducts (Table 6-3), i.e., (1) oxygen-containing copper (tough-pitch copper), (2) oxygen-free copper (deoxidized copper), or (3) oxygen-free copper (not deoxidized copper).

6.2.2.1 Tough-Pitch Oxygen-Containing Copper Tough-pitch copper normally has an oxygen content of 0.02-0.04%. This amount has a very low influence on the electrical, mechanical, and physical properties of pure copper. With oxygen copper forms the phase copper(i) oxide (Cu2O) (Fig. 6-4a and b), and has a eutectic (a-Cu + Cu 2 O), which precipitates and surrounds the ocsolid-solution copper crystals at the grain boundaries. As oxygen oxidizes the other impurities in the material, it can diminish their bad influence on the conductivity and cold working. Therefore in most cases in electrotechnology, oxygen-containing copper (E-Cu) is used and must meet a minimum electrical conductivity requirement (Table 6-3). Semiproducts of E-Cu 57 and E-Cu 58 have a content of 0.005-0.040% oxygen and an electrical conductivity of more than 57 and 58 m/Q mm 2 , respectively. When using oxygen-containing copper, the use of soldering, welding, or annealing in a

Table 6-3. Different,unalloyed copper materials3. Designation15 DIN 1787

Material number

ISO R1337

Oxygen-containing copper: E-Cu58 Cu-ETPC 2.0065 Cu-FRHCC E-Cu57 Cu-FRHC 2.0060 Cu-ETP

Composition range (wt.%)

Electrical conductivity ^(m/Qmm 2 )

Thermal conductivity X (W/mK)

99.90 O 0.005-0.040 >Cu 99.90 O 0.005-0.040

>58

386

> 57

386

> 58

386

~55

<38O

Special remarks

electrolytic tough-pitch (ETP) copper fire-refined high-conductivity (FRHC) copper

Main applications

for high conductivity: electrical engineering

Oxygen-free copper, not deoxidized, highest conductivity: OF-Cu

Cu-OF OFHC-Cu

2.0040

>Cu 99.95

free of deoxidizing elements, oxygen-free copper, oxygen-free, high-conductivity copper

Oxygen-free copper, deoxidized with phosphorus: SE-Cua

-

2.0070

SW-Cu

Cu-DLPC

2.0076

SF-Cu a

Cu-DHP

C

2.0090

>Cu 99.90 P 0.003 >Cu 99.90 P 0.005-0.014 >Cu 99.90 P 0.015-0.040

for the highest requirements: vacuum technology, electronics

oxygen-free, electrolytic copper

for low conductivity requirements: • electronics, material cladding

phosphorus deoxidized copper, low residual phosporus phosphorus deoxidized copper, high residual phosphorus

• chemical apparatus, construction engineering • pipelines, chemical apparatus, civil construction

k> or

35-53

240-360

DIN 1787, Bargel and Schulze (1994); b E-Cu - electrolytic copper (oxygen-containing); OF-Cu - oxygen-free copper (electrolytic); OFHC-Cu - oxygen-free high-conductivity copper; SE-Cu - oxygen-free electrolytic copper; SF-Cu - phosphorus-deoxidized copper (high residual phosphorus); SW-Cu - phosphorusdeoxidized copper (low residual phosphorus); c see Table 6-4.

"3 O

o 3 Q.

2
I O

CD

oo CO

290

6 Copper-Based Alloys

duce the electrical and thermal conductivity, especially phosphorus. However, phosphorus is used most frequently for deoxidizing (very brittle pre-alloy of CuPIO, DIN 17657). If there is excess phosphorus it will picked up in the copper solid solution. SW-Cu and SF-Cu are deoxidized coppers which are oxygen-free and contain a small amount of phosphorus; they are not required for their conductivity but other characteristic properties. They are used in chemical apparatus, pipelines, civil construction and engineering. SF-Cu can be welded and soldered.

1150

1100

•0.01 1050

(a) 1000 0

2

4

Cu

8

10

L

1100 1000

6

Cu2O (%)

cc

SL 900

6.2.2.3 Oxygen-Free Copper with High Conductivity

/ / a+Cu 2 O

|

800

|

700

/

Ol 1—

600

0 Cu

(b) 0.004 0.008 Oxygen (wt.%)

0.012

Figure 6-4. (a) Phase diagram for copper and copper oxide Cu 2 O (Bargel and Schulze, 1994; Dies, 1967). (b) Copper corner of the binary phase diagram for copper and oxygen (Dies, 1967).

reducing atmosphere can cause problems as a result of hydrogen-induced brittleness. 6.2.2.2 Deoxidized Copper

When using pyrometallurgical refining, the oxygen cannot be completely removed from the metal. The processing of oxygenfree copper is not done by electrolysis alone, and deoxidizing also adds elements, e.g., P, Si, Li, Mg, B and Ca. They combine with the oxygen to form oxides which remain in the slime. Surplus deoxidizing elements re-

The highest electrical conductivity and the best oxygen-free condition are found for copper (SE-Cu, OF-Cu) that is produced from good cathodes, remelted under a reducing atmosphere, and afterwards cast in ingot molds or extruded with a protective gas. This electrolytic copper is free from gassing elements and is used in vacuum technology. Some oxygen-free copper is produced in a vacuum and has no contact with oxygen from the melting of the cathodes to the final cooling; for that reason it contains less than 0.001 % oxygen (Table 6-3, Figure 6-4b). OFHC (oxygen-free high-conductivity) copper is produced by electrolysis and remelted and cast under a protective gas atmosphere or in vacuum, therefore it has an electrical conductivity of more than 58.5 m/Q mm 2 . It is resistant against hydrogen enbrittlement, has an oxygen content of less than 0.001%, can easily be welded and soldered, and no deoxidizing elements are added. A special version is SE-Cu; this can also be produced after electrolysis by additions of deoxidizing elements, e.g., lithium, car-

6.2 Technology and Requirements of Copper

291

Table 6-4. National and international standards and designations for pure coppera. Germany DIN 1708 DIN 1787

France NFA 53-100

Great Britain BS 2872 etc.

Italy UNI 5649

Sweden SIS

Switzerland VSM 10826

KE-Cu El-Cu58 E2-Cu58 E-Cu57

Cu/d Cu/al Cu/a2 -

_ C101 C102 -

Cu-CATH Cu-ETP -

_ 145010 145013 -

F-Cu OF-Cu b -

Cu/a3 Cu/c2 Cu/cl _ _ Cu/b

C104 C103 -

Cu-FRTP Cu-OF b _ Cu-DLP Cu-DHP

145013 145011 _

SE-Cub SW-Cub SF-Cub

_ C106

145015

Austria ONORM M3401

U.S.A. ASTM B224

ISO R1337

_ Cu-ETP _ Cu-OF b Cu-OFE

Cu-K Cu-E Cu-F -

CATH ETP FRHC FRHC

Cu-CATH Cu-ETP Cu-FRHC Cu-FRHC

Cu-C Cu-OF b -

FRTP OF b OFE

Cu-FRTP Cu-OF b -

Cu-OFLP — Cu-DHP

Cu-SEb Cu-DLP Cu-DHP

— DLP DHP

Cu-PHC Cu-DLP Cu-DHP

a

DKI (1982); CATH, K, C - cathode copper; KE-Cu - cathode electrolytie copper; E-Cu - electrolytic copper; ETP - electrolytic tough-pitch copper; FRHC - fire-refined high-conductivity copper; FRTP - fire-refmed tough-pitch copper; F-Cu - fire-refined copper; OFE - oxygen-free electrolytic copper; PHC - phosphorusdeoxidized high-conductivity copper; DLP - deoxidized low phosphorus residue; DHP - deoxidized high phosphorus residue; OFLP - oxygen-free low phosphorus residue; b see Table 6.3.

bon, and a small amount of phosphorus. It therefore contains less than 0.003 % P. CuPHC (phosphorus-deoxidized high-conductivity) and Cu-PHCE have the same properties. 6.2.3 Copper Standards

Designations, composition ranges, required properties, and terms of delivery of copper are compared with national and international (ISO) standards and listed in Table 6-4. The production process, castings, semiproducts, and the conductivity are mostly taken into consideration in the classification of copper. The classification of pure, unalloyed copper materials depends on their purity, as determined by the production process (Table 6-3). A special standard is OFHCcopper (oxygen-free high-conductivity) with 100% I ACS conductivity equal to 58.0 m/Q mm 2 at 20 °C (IACS: International Annealed Copper Standard). Other

alloys are compared with this pure copper standard, which is soft annealed between 400 and 500 °C, and has a density of 8.89 g/ cm3 (DKI, 1982). Copper standards are used to ensure the identification of specific chemical compositions of wrought and cast copper alloys. A six-digit UNS number (Unified Numbering System) is used by the ASTM (American Society for Testing and Materials) in the United States and Canada. Other national standards are also applied worldwide, e.g., BS (British Standard) in the United Kingdom or DIN (Deutsches Institut fur Normung) in Germany. 6.2.3.1 Deutsches Institut fur Normung (DIN) Standards

The DIN standard distinguishes different alloying classes, and depends on the alloying elements and whether the products are wrought, cast, or semiproducts (DIN, 1991). It describes the composition

292

6 Copper-Based Alloys

range of alloys with the DIN designation, which determines the different alloying elements, followed by the content in weight% (mass%) of the main designated elements, e.g. Cu-A110-Fe3-Mn2. 6.2.3.2 American Society for Testing and Materials (ASTM) Standards Copper and its alloys are divided according to two categories: UNS numbers C10000-C79 999 for wrought copper alloys, and C80 000-C99 999 for cast copper alloys (ASTM, 1991). Within these two categories, the compositions are grouped

into families of copper and copper alloys, as shown in Table 6-5 (CDA, 1990; Covingtonetal., 1992). 6.2.4 Copper Consumption The largest quantity of copper is consumed by the electrical industry (50 %) for cables, wires, and products with high electrical conductivity (Metallgesellschaft, 1993). The civil and construction industry is the second most important consumer using copper and copper alloys in heating, water supplying, sanitary technology, and also for roof coverings, gutters, and wall

Table 6-5. ASTM system of categorizing copper and its alloysa. Copper family Coppers

Brasses

Bronzes

Copper-nickels Nickel-silvers

UNS numbers CIO 000--C15999 C80000--C81299

Characteristics wrought and cast copper with minimum content of 99.3 % Cu

^) high-copper alloys for wrought products, less than 99.3 % but C16000--C19999 Lmore than 96 % Cu, which do not fall into any other alloy group; C81300--C82999 f for cast alloys in excess of 94 % Cu, to which silver may be added * for special properties C20 500--C28 599 Cu-Zn alloys j with or without other designated C31 200--C38 599 Cu-Zn-Pb alloys (leaded brasses) \ alloying elements, such as iron, J aluminum, nickel, and silicon C40400--C49999 Cu-Zn-Sn alloys (tin brasses) C83 300--C85 899 Cu-Sn-Zn cast alloys (red, semired, and yellow brasses) C86100--C86399 Cu-Sn-Zn-Mn alloys (manganese bronzes) C86400--C86899 Cu-Sn-Zn-Mn-Pb alloys (leaded manganese bronzes) C87 300--C87999 Cu-Zn-Si cast alloys (silicon brasses and bronzes) C50100--C52499 Cu-Sn-P wrought alloys (phosphorus bronzes) C53 200--C54899 Cu-Sn-Pb-P alloys (leaded phosphorus bronzes) C60 600--C64499 Cu-Al alloys (aluminum bronzes) C64700--C66199 Cu-Si alloys (silicon bronzes) C90200--C91 799 Cu-Sn cast alloys (tin bronzes) C92200--C94599 Cu-Sn-Pb alloys (leaded tin bronzes) C94700--C94999 Cu-Sn-Ni alloys (nickel-tin bronzes) C95 200--C95 900 Cu-Al alloys (aluminum bronzes) C70100--C71999 \ Cu-Ni alloys with or without other elements C96200--C96999

C73 000--C79999 1 rCu-Ni-Zn alloys, contain no silver C97 300C98 200- C98 999 Cu-Pb cast alloys, without tin or zinc Leaded coppers Special copper alloys C99 300- C99999 chemical compositions do not fall into any of the above categories ASTM (1991).

6.2 Technology and Requirements of Copper

1979

1991

293

Electrical industry Consumer and lighting products

Figure 6-5. Application fields for copper (in %) in the western world from 1979 to 1991 (Metallgesellschaft, 1993).

claddings. The automotive industry uses copper-containing materials for radiators, heat exchangers, and wiring systems, The development of the total consumption between 1979 and 1991 is classified according to the different industrial fields in Fig. 6-5. Some users try to replace the expensive copper material by other materials, e.g., aluminum for power cables and transmission lines, and for radiators in the automotive industry, or optical fiber cables in the telecommunication and computer industries because of their higher efficiency. However, copper finds new applications, e.g., through the electrification of many areas in the household, and via telephones, videos, or automobiles. The use or replacement of copper depends on two considerations, the saving of energy and the cost of the raw copper material. The worldwide mine output of copper in the main producer countries grew between the years 1980 and 1991 by 17.4% from 7.7 million tonnes (Mt) to 9.1 Mt (Table 6-6 a). In the same years the output of refined copper grew from 9.3 to 10.6 Mt, i.e., by about 15% (Table 6-6 b). The consumption of refined copper in the main consumer countries has experienced a growth of 12.8% from 9.5 to 10.7 Mt worldwide

(Table 6-6 c). So copper is an important material, which is always needed. The per capita consumption in the individual countries has shifted from 1980 to 1990. While the East European countries experienced reduction, e.g., the Soviet Union (CIS) from 4.9 to 3.4 kg/person, the amount used in the newly industrialized countries increased greatly, e.g., Taiwan increased its consumption by more than three times to 14.7 kg/person in 1990; this is higher than that of Germay (14.2), Japan (12.8), and the U.S.A. (8.6 kg/person) (Metallgesellschaft, 1993). The distribution of copper castings and copper alloys in Germany was nearly the same in the years 1970 and 1980 (Table 6-7). Nearly half this amount represents Cu-Zn castings, and to a much lower extent unalloyed copper castings (Weisner etal., 1982). The price of copper at about 1 US dollar/pound ( ~ 4 DM/kg) certainly makes it one of the more valuable nonferrous metals - the same goes for brass and tin bronze. Therefore the recycling of copper alloys is important, because scrap is an additional source of raw material, the production of secondary copper materials is less expensive, copper production from

294

6 Copper-Based Alloys

Table 6-6. Copper production and consumption by the main countries worldwide in the years 1980 and 1991 (Metallgesellschaft, 1993).

Alloying group

(a) Mine output in 1000 tonnes. Country

1980

1991

U.S.A. Chile C.I.S. (U.S.S.R.) Canada Zambia Zaire Peru Poland Philippines Australia

1181 1106 980 716 596 505 367 343 304 244

1808 Chile 1635 U.S.A. C.I.S. (U.S.S.R.) 870 774 Canada 423 Zambia 381 Peru 370 China 330 Poland 315 Australia 235 Zaire

World total a

7700

Copper-zinc castings (brass) Copper-tin-zinc-(lead) castings Copper-tin casting alloys Copper-lead, Copper-aluminum, and other casting bronzes Unalloyed copper castings Copper castings with up to 10% nickel Total a

World total

9100

Country

1980

U.S.A. C.I.S. (U.S.S.R.) Japan Chile Zambia Canada Germany Poland China Belgium/ Luxembourg

1686 1300 1014 811 607 505 469 357 295

World total

9300

289

1991 U.S.A. Chile C.I.S. (U.S.S.R.) Japan Canada Germany China Zambia Poland Belgien/ Luxembourg World total

2001 1238 1200 1076 538 523 500 390 378 286 10 600

Growth = 15%.

(c) Refined copper consumption in 1000 tonnes. Country

1980

1991

U.S.A. C.I.S. (U.S.S.R.) Japan Germany France United Kingdom Italy China Belgium/ Luxembourg Brazil

1868 1300 1153 888 433 409 388 370

U.S.A. 2058 Japan 1613 Germany 1005 C.I.S. (U.S.S.R.) 920 China 570 France 481 Italy 470 Taiwan 397 Belgium/ Luxembourg 372 South Korea 343

World total

9500

a

Percentage of total production 1970

1980

45 29 16

45 29 11

5 3

4 4

2

7

100

100

Weisner et al. (1982).

Growth =17.4%.

(b) Output of refined copper in 1000 tonnes.

a

Table 6-7. Consumption of copper castings and copper alloys in Germany in the years 1970 and 1980a.

Growth=12.8%.

304 246

World total

10 700 a

scrap requires less energy, the recovery conserves the deposits of raw materials and energy, and the need for landfill space is reduced, reducing the cost of storing wastes and residues. So more of the growing recycling potential of up to 70 % (Fig. 6-6) must be used for copper applications, at present about 10 Mt/year worldwide. The aim must be to have a residuefree production and consumption cycle. 6.2.5 Requirements with Regard to Copper and Its Alloys

As we have seen, using different fabrication processes, copper is produced from its oxidic or sulfuric ores and then offered to the market in various qualities, with special purity, and in different delivery forms (Table 6-8). The applications of copper are determined by basic considerations as follows: • Copper is the second best conducting metal after silver and has a specific electrical conductivity of about 58 x 106 S/ m (E-Cu). Because of the high electrical conductivity of copper, it is necessary to

6.3 Properties and Behavior of Copper

H10000

i70

8000 - Copper consumption

60.S

6000

50

4000

40

2000

30

1970

295

1975

1980 Year

1985

produce and apply copper in as pure a condition as possible. Pure metals are relatively soft and less strong. They can be strengthened by alloying and then by solid-solution hardening, precipitation hardening, grainboundary hardening, dispersed particles, phase transformation, etc. Hence the need to alloy pure, soft copper, to insert particles (dispersion hardening), or to alter the microstructure by thermal and mechanical treatments. During plastic (cold) deformation, dislocations are produced. A high density of dislocations leads to higher strengths. Therefore, copper and copper alloys for applications requiring high strength (e.g., for long conductor lengths) need to be in a hard, deformed semiproduct condition. Using mechanical and thermal fixation and connection processes (screwing, bolting, clamping, welding, soldering) to form end products, the copper material is fastened or heated excessively locally. As a result of this, the strength increase can be partially or totally canceled, i.e., the material becomes softer. Therefore copper and copper alloys must be strengthened using special alloying elements and special heat-treatments.

1990

20

Figure 6-6. Development of copper consumption (in 1000 tonnes) and recycling potential (in %) worldwide in the years from 1970 to 1990 (Metallgesellschaft, 1993).

• The chemical resistance of copper and bronze has been known since ancient times (bronze age). Copper forms, on oxidation, a tight, strongly fixed oxidation layer, the green patina. Copper materials are therefore corrosion resistant and used in, e.g., chemical plants or seawater stations. From these considerations and qualified requirements concerning copper materials, the most suitable or favorable material and its condition are chosen for the specific application (Table 6-9). The many favorable properties, the various applications, and the great demand entail a wide range of products and preserve the ancient copper metal as modern and attractive as before.

6.3 Properties and Behavior of Copper The main properties of copper are summarized as follows: • Copper is corrosion resistant - it is resistant for centuries against weather and wind; in seawater and corrosive liquids copper and low-alloyed copper are resistant over a long period of time: ap-

296

6 Copper-Based Alloys

Table 6-8. General outline of principal applications of unalloyed copper and the main copper alloys. Decisive property High electrical conductivity

High thermal conductivity

Product form

Examples of application fields cables for telephone, transmission lines, deep-sea cables, telecommunication and signalization circuits, accessory parts for radio and television winding of electrical machines, generators, motors, transformers, instruments, superconductors conducting lines in buildings and vehicles, overground and underground cables, conducting rails, battery connections, clamp profiles, switch accessories wires for rails and chains, connections of rail parts, conductors for electric railways, tramways and buses electrodes and conductors for welding apparatus, spark erosion, melting furnace, etc., lightning-rods, earthing systems, power utilities

copper wire

heat exchangers, heat transmitter for cooling water, oil, liquids, heat transfer for solar plants and for heat recycling tubes, vessels, and spirals for heating and cooling plants, air-conditioning, hot-water vessels in buildings, apparatus for distillation, condensation, and cooking, chemical and food apparatus, refrigeration die casting and plates for casting metals, foundries

copper and alloyed tube, plate

water-cooled jets and accessories, cooling forms of blast furnace High resistance against corrosion

Easy manufacturing, good junction

apparatus, tubes, vessels in chemical, food and brewery industries roofs of buildings, gutters, lightning-rods, insulation against humidity gas and water pipes, outside electrical conductors, sanitary installation, water boilers, store vessels, heating tubes construction parts in all industries, appliances semiproducts like plates, tubes, rods, profiles, strips, nails, rivets containers, boxes, deep-drawn parts, pipes, valves, fittings arts and crafts, plastics, coins

plications in civil construction protect against corrosion, e.g., roofs or installations. Copper is electrically and heat conductive - a computer, an automobile, a generator, or a telephone will not work

copper wire copper and brass strip, copper wire

copper and low-alloyed copper wire, profile, strip copper wire, bar, strip

copper tube, plate, brass and alloyed tube, plate copper and alloyed strip, plate copper rod, tube copper and alloyed plate, tube copper plate copper and brass tube, rod all all copper and brass plate, tube, casting copper and alloyed copper

without copper for modern technology; copper and low-alloyed copper are the most important conductors of electricity and heat. Copper is healthy - so copper as an element is important for human life; as a

6.3 Properties and Behavior of Copper

297

Table 6-9. Comparison of electrical and mechanical properties of some selected copper alloys. Designation

Alloying constituents (wt.%)

Cu-Sn2

Snl.5-2.5, Zn0.3

>30

Cu-Sn6

Sn5.5-7.5, ZnO3

10-20

G-Cu-SnlO-Zna G-Cu-Sn5-Zn-Pba

Sn9-ll, Zn4 Sn5-6.5, Zn4-6

7 9

Cu-Cdl Cu-Mg0.7 Cu-Crl

CdO.9-1.3 MgO.5-0.8 CrO.3-1.2

36-48 <36 «50

Cu-Si2-Mn Cu-Ni20-Mnl0 Cu-Mnl2-Ni

Sil.5-3.5

4-6 1.5 1.5

G-Cu-Pb22 G-Cu-Pb5-Sn

Pbl8-23,Ni,Fe, Sn Pb4-6, Sn9-ll

Cu-A15 Cu-AllO-Ni

A14-6, Ni0.8 Ni3-6, Fe2.5-5

8 6

Cu-Zn44-Pb2 Cu-Zn33

additions 1-2.5

18 16

Cu-NilO-Fel-Mn Cu-Nil8-Zn20 Cu-Ni25

NilO.4, Fel.5, Mn0.8 Nil8, Zn20 Ni25

a

Mnl2-15

Electrical conductivity x (m/Q mm2)

5.3 3.5 3.1

Tensile strength Rm (N/mm2) 270 480 370 660

soft-heat-treated rolled soft-heat-treated rolled

Applications

contacts, screws conductive springs, strips, bands

250- 300 hard 200- 240

sliding bearings good castings

400- 600 600-1000 400- 600

conductors conducting pieces conducting rails

> 850 spring hard 350 soft 330

springs resistors resistors

160 240

gliding planes and bearings

350 560

seawater equipment

500 360 soft

hard brass construction brass for soldering

420-450 400-610 >290

seawater, offshore cutlery, fittings coins, claddings

G = sand-mold casting.

trace element in different foods it has an effect, e.g., in forming pigments, reducing cholesterol, and protecting against dental caries. Copper is beautiful - at all times during human existence, copper has been present in arts, for decoration and jewellery; this is based on its colorful and splendid appearance and surfaces, which its alloys also possess. Copper is easy to deform - because of its crystal structure, copper can be formed, rolled, drawn, plated, or extruded; in mechanical and chemical technology, different types, shapes, or surfaces are used in the form of plates, sheets, pipes, wires, or claddings.

6.3.1 Properties of Unalloyed Copper The most important technical characteristic of copper is its electrical conductivity. Pure, unalloyed copper has, after silver, the highest electrical and thermal conductivities, which depend to a large extent on its purity and are reduced by alloying (Table 6-3, Fig. 6-7). If foreign atoms are incorporated in the structure of copper, they can substitute the copper atoms (substituted structure) or sit between the copper atoms (intermediate structure). There they produce an interference potential, which disturbs the electrical field and the movement of electrons through the matrix. Thus the movement of

298

^

6 Copper-Based Alloys

0.016

60

0.018 -

55

E

a . 50

^0.020 .3 0.022

c o

•E 0.024

h I 40 LU

0.026

35

0.028L

30

0.02

0.04 0.06 0.08 Additions (wf.%).

0.10

0.12

Figure 6-7. Influence of impurities on the electrical conductivity x and specific electrical resistance ge of pure copper (Pawlek and Reichel, 1956).

electrons is impeded, and the electrical and thermal conductivity are impaired. The conductivity can be reduced noticeably by even small contaminations of only 0.01 %, in particular of phosphorus, iron, or cobalt. Impurities that are dissolved in the solid condition in copper (mixed crystals) reduce the conductivity, even as small additions. In particular, phosphorus, as a good deoxidizing element, impairs the conductivity considerably. On the other hand, elements that are not dissolved in the solid solution have virutally no influence as small additions on the conductivity, e.g., oxygen from the production process or lead and tellurium which sit at the grain boundaries. The values of the main properties of highly conducting copper (E-Cu) are given in Tables 6-1 and 6-10 (DKI, 1982; Wallbaum, 1964) (UNS numbers CIO000C15 999). Because of its high thermal conductivity, copper is used in heat exchangers of combustion engines, in coolers and radia-

tors, in heating coils, in chemical plants, and in breweries. The thermal and electrical conductivities are proportional to one another. For heat transfer in heat exchangers, the formation of an oxidized layer (corrosion protection) is more important and decreases the conductivity more than the impurities in the copper. The resistance to corrosion in neutral or alkaline solutions is also of importance (water pipes). The formation of oxide layers or crusts must be avoided. In contact with air, a protective layer of basic copper carbonate (green patina) is formed with carbon dioxide. As an end product, unalloyed copper is used for many different applications in the electrical industry. In oxidizing acids (e.g., HNO 3 ) copper is corroded. In a deoxidizing, hydrogencontaining atmosphere above 500 °C, the oxygen-containing copper can be cracked by hydrogen Cu 2 O + H 2 -> 2 Cu + HLO

(6-1)

The hot-water steam cracks the copper along the grain boundaries (formation of bubbles, welding cracks, etc.). At temperatures above 1000°C oxygen diffuses into copper along the grain boundaries. The technological properties are based on the good deformation in hot and cold conditions of the face-centered cubic (f.c.c.) crystal structure of the oc-copper (Fig. 6-39). Unalloyed copper has a low strength and a high fracture toughness, e.g., normalized copper: tensile strength i? m ^220 N/mm 2 , fatigue strength =70 N/ mm 2 , and fracture toughness at room temperature = 80-140 J/cm2. The strengthening of copper by cold deformation is not very high (Fig. 6-8), but a high reduction of the area of up to 90 % can be obtained. This cold deformation can be removed at a low annealing temperature (Fig. 6-9) by recovery and recrystal-

6.3 Properties and Behavior of Copper

299

60

I 40 Figure 6-8. Influence of cold deformation on the tensile strength, Brinell hardness, and fracture elongation of electrolytic copper (Bargel and Schulze, 1994).

20

20

40 60 80 Reduction area (%)

lization. All of the strength-related properties depend on the content of dissolved impurities, which block the dislocations. Hence unalloyed and annealed copper has a low strength. With increasing temperature, the strength of copper decreases (Fig. 6-10) and the electrical resistance increases (Fig. 6-11). Welding copper is difficult because of its high thermal conductivity. The heat used for melting the connecting parts is conducted to a large extent into the material being welded. So with arc or autogenous welding, only oxygen-free copper can be used to avoid copper oxide formation or the diffusion of hydrogen or nitrogen, which would lead to brittleness of the material (Steffens and Sievers, 1990).

500 E

100

The microstructure of SF-Cu has elongated grains in the deformed condition which can be released by annealing to coarse grains (Fig. 6-12). 6.3.2 Low-Alloyed Copper

With low-alloyed copper (up to about 5 % alloying elements), the electrical conductivity is reduced slightly, but the strength of the copper can be increased considerably through small quantities of alloying elements. This can be achieved by solid-solution hardening (silver, arsenic) or by precipitation hardening (chromium, zirconium, cadmium, iron, or phosphorus). The formation of mixed crystals is unfavorable for the electrical and thermal

Annealing time, f = 1h 200 - 9 5 % cold deformation 150

400 -

% cold deformed soft

* 100 300

200

10%

100 200 300 400 Annealing temperature (°C)

50

500

Figure 6-9. Recrystallization of cold-deformed copper by heat-treatment.

£

0

100 200 Testing temperature (°C)

300

Figure 6-10. Stress rupture of SE-copper after 1000 h depending on the temperature and amount of deformation (VoBkiihler, 1955).

300 * r 400 e

6 Copper-Based Alloys

r

e c: 0.04

"^200
wo

v 100

c cu

•- o

lee trical resisi ivii

^300

LU

0.03

60

0.02

40

0.01

20 n

0

-200

-100

0 100 200 Temperature (°C)

300

conductivity, while the formation of precipitates results in an optimum combination of strength and conductivity (Grafen, 1991; DKI, 1976); the solid solution hardening is rather small and reduces the electrical and thermal conductivity.

Figure 6-12. Rod of SF-copper (99.94% Cu; 0.024 %P): (a) cold deformation of about 30 %, elongated grains; M = 500. (b) soft-annealed, a-grains with copper oxide (round inclusions).

0 3

Figure 6-11. Electrical resistivity Qe, electrical conductivity x, and tensile strength Rm of SE-copper depending on the temperature (Guillery et al., 1991; DKI, 1982).

400

Low-alloyed copper is used for special electrically conducting and engineering applications, e.g., commutator segments, welding electrodes, spring contacts, automats, etc. (UNS numbers C16000C19999). Cu-Cd, Cu-Zr, or Cu-Fe-P alloys are used in the electronics of semiconductors, and new materials have been developed for thin electronic strips e.g., Cu Ni2-Si 0.7 and Cu-Ni3-Sil (Ruchel, 1989). Automobile coolers are produced from Cu-Ag alloys, and hardened Cu-Be alloys with up to 2% Be can be used for springs or contacts. A small amount of tellurium (up to 0.5%) results in heterogeneous inclusions in the solid microstructure. This alloy can easily be machined by automatic cutting apparatus. Low-alloying elements for copper, which are favorable for the strength and conductivity (reduced by only 20%), include Ag, Cd, Cd-Sn, and Be. Those that are not very acceptable in terms of reduced conductivity (reduced by between 20 and 50%) are small amounts of O 2 , Sb, As, Si, Fe, and P (Fig. 6-7). For high strength, suitable low-alloying elements are Zr, Cr, Cr-Zr, and Be (precipitation hardening). In addition, the strength of these alloys can be increased by heat-treatment.

6.3 Properties and Behavior of Copper

301

Table 6-10. Physical and mechanicala properties of unalloyed and low-alloyed (nonhardenable and hardenable) copper b. Material

Linear thermal expansion a (25-300 °C) (10~6/K)

Electrical conductivity at 20 °C, x (m/Q mm2)

Thermal conductivity at 20 °C, X (W/m K)

Unalloyed: E-Cu58 E-Cu57 SE-Cu SF-Cu

17.7 17.7 17.7 17.6

>58 >57 57-58 45-52

394 386 386 300-350

200200200200-

360 360 360 370

45-120 45-120 70-120 50-105

Nonhardenable: Cu-AgO.l Cu-Cd0.7 Cu-Mg0.4 Cu-Fe2-P

17 17 17.6 16.3

55-57 53-56 >36 17-38

385 355-370 240 200-260

250340400300-

360 380 600 520

70-120 100-120 110-175 75-145

Hardenable: Cu-Bel.7 Cu-Co-Be Cu-Cr-Zr Cu-Ni2-Si

17 18 17 16

8-18 25-37 26-48 10-23

92-125 192-239 170-320 67-120

390-1380 250- 950 370- 470 260- 640

90-420 70-260 120-170 70-190

a

The mechanical properties depend on the thermo-mechanical treatment; (1992).

6.3.2.1 Nonhardening Copper Alloys Nonhardening copper alloys contain alloying elements in solid solution; these alloying elements, e.g., Ag, As, Cd, Mg, Mn, Si, S, and Te, do not generally increase the mechanical properties or decrease the conductivity very much, but have special, favorable properties (Table 6-10). In small additions, silver increases the softening temperature after quenching (Fig. 6-13), and therefore the strength and the creep resistance at elevated temperatures. Arsenic increases the softening of cold-worked copper alloys and the scale from gases, but decreases their high conductivity. Cadmium also increases the static and the dynamic strength at elevated temperatures, and the wear resistance. Magnesium has virtually the same effect. Manganese increases the basic strength

Tensile strength,

Brinell hardness, HB

(N/mm2)

b

DIN 17666, Arpaci and Bode

and corrosion resistance of copper; like silicon and manganese, tellurium and sulfur improve the machinability without greatly influencing the conductivity. With all these low-alloyed coppers, strengthening is only possible by cold deformation. 6.3.2.2 Hardening Alloys Hardening alloys, in contrast to nonhardening alloys, improve the strength by heat-treatment (solution treating) and precipitation hardening using, e.g., Be, Cr, Ni, and Si (Table 6-10). In this case, after solution treating (homogenizing) in the ot-phase, precipitates are formed by annealing or tempering in the binary phase field a + X (where X = intermetallic phase) (Fig. 6-14) (Fischer, 1978). These particles block microstructural and dislocation sliding and so increase the strength, hardness,

6 Copper-Based Alloys

Figure 6-13. Modification of softening temperature of OF-copper by alloying elements (DKI, 1976). Preparation: intermediate annealing 0.5 h at 600 °C, cold deformation with 75 % area reduction. 0.02 0.03 0.04 Additions (wt.%)

and other mechanical properties, while reducing the electrical and thermal conductivities. 6.3.3 Conductive Bronze

Low-alloyed copper alloys for electrically conducting applications are described

1400 -

0.05

by their electrical conductivity and not by the alloy content. They are grouped according to the collective name of conductive bronze (Table 6-11). Because of their higher strength due to the alloying elements, and because of their improved wear resistance together with the best possible electrical conductivity (Fig. 6-15), they are used for outdoor and overhead lines (ASM, 1979; DKI, 1976).

6.4 Properties and Behavior of Copper Alloys

0.2

0.4 Cr (wt.%)

0.6

0.8

Figure 6-14. Precipitation hardening of Cu-Cr alloys (Fischer, 1978). X = intermetallic phase; 1= homogenizing; 2 = temper annealing.

Beside the technical application of the largest amount of unalloyed copper and low-alloyed copper as electrical conductors, many copper alloys are produced mainly as materials with elevated strength, as cheaper materials than copper, and as materials with special or characteristic properties. Many copper alloys can be produced or manufactured by deformation (rolling, drawing, extruding) or as cast ma-

6.4 Properties and Behavior of Copper Alloys

303

Table 6-11. Properties and applications of conductive bronzes with minimum electrical conductivity values3. Conducting bronze

Electrical conductivity,

Tensile strength,

(m/O mm2)

(N/nun2)

I (L45)b (min. 99% Cu)

>48

II (L35) (min. 98% Cu) III (LI 5) (min. 96% Cu) a

DKI (1976);

b

Composition range

Applications

500-520

Cu-CdO.5-1 Cu-Mg0.1

conductors, contacts, electrode supports, short-circuit rings

>36

560-680

Cu-Mg0.3-0.5 Cu-Cdl-Snl

electrodes, high-voltage switches, cylinder heads

>18

660-740

Cu-MgO.5-0.7 Cu-Sn2-3 Cu-Snl.2-Znl.2

fittings, seals, cooling boxes, armatures

L = electrical conductivity value.

terial (sand casting, die casting, pressure casting). Their properties are generally influenced and changed by three processes: alloying, heat-treatment, and deformation, The properties of copper and the main copper alloys have wide ranges depending

0 Cu

1 2 Increase of s t r e n g t h , /? p o.oi. through alloying additions (wt.%)

3

Figure 6-15. Influence of alloying additions (wt.%) (amounts on curves) on the strength Rp0 0 1 and the electrical conductivity x of copper (DKI, 1976).

on the treatment. They are summarized in Table 6-12. The relation between the electrical conductivity and the Brinell hardness (parallel to the strength) for some selected copper alloys is shown in Fig. 6-16. Above the median line is the field of hardened, low-alloyed copper alloys, which are suitable for applications requiring high strength and high conductivity. Below the median line is the field of age-hardenable copper alloys, the hardness and conductivity of which can be improved at the same time by the hardening treatment. Some of the copper alloys have historical names, like brass (Cu-Zn), bronze (CuSn), red bronze (Cu-Zn-Sn), nickel-silver (Cu-Ni-Zn), and tombac (from Cu-Zn5 to Cu-Zn33). In addition, special alloys with special properties are generally alloyed with two or more elements (multicomponent alloys). 6.4.1 Copper-Zinc Alloys (Brasses) The Cu-Zn alloys (brass or special brass) are the most important copper alloys. Brasses are alloys with the main elements copper and zinc and contain, for

304

6 Copper-Based Alloys

Table 6-12. Comparison of the physical and mechanical properties of copper and the main copper alloys (depending on treatment). Material

(g/cm3)

Melting temperature, solidification range (°C)

Temperature of hot deformation (°C)

Modulus of elasticity E (kN/mm2)

Shear modulus G (kN/mm2)

Specific heat capacity cp 20-100°C (J/gK)

8.93 8.3-8.8 8.8 8.5-8.7 8.2-8.7 7.5-8.2 8.9

1083 895-1045 910-1040 950-1100 865-1030 1030-1080 1060-1290

800-950 700-850 700-800 600-950 600-900 600-950 620-900

100-130 104-124 112-128 125-140 130-133 105-127 130-165

46.4 40-46 42-43 51-52 43-45 -

0.38-0.45 0.39 0.37 0.42 0.45 0.38-0.41

Electrical conductivity at 20 °C, x (m/Q mm2)

Tensile strength,

Fracture elongation,

Brinell hardness, HB

(N/mm2)

(%)

35-58 15-33 7-12 3- 5 13-27 5-10 2- 6

200- 360 240- 610 330- 630 370- 610 390-1500 370- 730 400- 570

2-45 12-50 6-60 13-45 2-35 5-35 14-35

Density Q

Cu Cu-Zn (brass) Cu-Sn (bronze) Cu-Ni-Zn Cu-Be Cu-Al Cu-Ni Material

Linear thermal Thermal expansion, a conductivity 20-100°C at20°C, X (W/m K) (10~6/K)

Cu Cu-Zn (brass) Cu-Sn (bronze) Cu-Ni-Sn Cu-Be Cu-Al Cu-Ni

16.8-17.2 18.2-20.3 18.2-18.5 16.5-19.5 17.0-17.6 17 -18 16 -17.5

240-385 117-240 65- 90 2 3 - 35 113-239 50- 83 30-130

55-110 50-190 65-200 85-185 70-400 90-210 85-190

oCuBe2

350

o CuBe1.7 300 o CuBeU 250

200

CuZn23Al6Mn4Fe3 + \ o CuCoBe

CuNi 2Si v CuSi3Mn \ o + \ CuNi1.5Si , CuZn40Al2 150 l h : . oi1" + +CuZn39Pb2 L lr u7znn1i i b ii+ +CuMg0.7 n + CuZn35Ni2+ CuZn40 •CuZn40Mn2 100

CuM

3°-4SF-CuTecuAg + + + + E-CuTe

Se-Cu

50

10

20 30 40 50 Electrical conductivity, x (m/&mm 2 )

60

Figure 6-16. General overview of copper alloys. Relation between Brinell hardness, HB10 (strength, Rm) and electrical conductivity, x (DKI, 1976). O = hardened material; + = hard-drawn material.

6.4 Properties and Behavior of Copper Alloys

technical purposes, at least 50% Cu. Besides this they contain up to 3 % Pb and additions of other elements, e.g., aluminum, iron, manganese, and tin (ASM, 1979; DKI, 1966, 1991a; Beitz and Kiittner, 1987; Wieland, 1986). Three groups can be distinguished: Cu-Zn alloys, Cu-Zn alloys with lead, and Cu-Zn alloys with other alloying elements. Zinc is added to increase the strength and make the material cheaper, but it also reduces the melting temperature, the possibility of deformation, and the electrical conductivity. Cu-Zn alloys with more than 67% Cu used to be known as tombac. The higher strength of these alloys in comparison to copper is due to solid-solution hardening, which causes the strength and hardness to increase (Figs. 6-17 and 6-18). Brass is standardized for deformable

3

80 -

20 30 Zn (wr.%)

40

Figure 6-17. Mechanical properties of Cu-Zn alloys measured at 20 °C in the soft-annealed condition (DKI, 1991a).

305

20°C 200°C

20 30 Zn (wt.%)

Figure 6-18. Tensile strength, Rm, of Cu-Zn alloys versus the Zn content at different temperatures (DKI, 1991a).

alloys (DIN 17 660; UNS numbers C20 500-C49 999) and for cast alloys (DIN 1709; UNS numbers C83 300C87 999) (DIN, 1991; Covington et aL, 1992). According to the Cu-Zn phase diagram (Fig. 6-19), two types of alloy can be differentiated (a and a + (3). Cu-Zn alloys are single-phase, homogeneous alloys with an amixed crystal up to ~ 3 7 % Zn (Fig. 6-39); the zinc is completely dissolved in the copper matrix. These oc-alloys are good for noncutting forming, in particular for all deformation processes, because of their face-centered cubic (f.c.c.) structure with easy dislocation glide (soft brass). The strength increases with a higher zinc content. The highest deformation grade contains 28% Zn; Cu-Zn28 is used for extreme deep drawing of, e.g., lipstick and pocket lighter sleeves, or ammunition. With a zinc content of greater than 37%, heterogeneous oc + P alloys are formed (Fig. 6-20). These are easy to cut mechanically and to hot deform, but are difficult to cold deform because of the hindrance to dislocation movement (hard brass). The (3-phase is body-centered cubic

306

6 Copper-Based Alloys

1100 •1083°C

Liquid melt+a

Liquid melt

1000

902°C

900 Hot defamation

_ 800

ood hot deformabiliry^

« 700

Good hot deform ability hot deformability

fe 600 Q.

500 400 300 200

0 Cu

10 CuZn10

20 / 30 CuZn20 / CuZn30 Good cold deformability, bad machining

40 \ 50 CuZn40 \ Zn(wt.%) Less good cold deformability, good machining

Figure 6-19. Copper-zinc phase diagram, showing Cu-Zn alloys (Hansen and Anderko, 1958) and regions of hot deformation, recrystallization, and stress-free annealing (DKI, 1991a).

(b.c.c.) and more brittle. This results in a rapid reduction in the impact strength with a simultaneous increase in the hardness and tenside strength. Lower impact strength and more heterogeneity are preferred for cutting manufacturing, e.g., drilling, turning, or milling to form short chips. For Cu-Zn44 the strength is at its maximum of about Rm = 540 N/mm 2 .

Figure 6-20. Micro structure of pressed profile CuZn39-Pb3; ct + p structure; fi solid solution dark.

The second group of Cu-Zn alloys includes lead-containing or cast brass for easy automatic cutting. These alloys contain either brittle particles of the |3-phase or precipitates of lead, neither of which are good for deformation. The best cutting alloy contains about 1-2% Pb. The lead addition (up to 3 % Pb) does not alloy, but takes the form of fine, spherical precipitations and so aids machining. Cu-Zn42-Pb is optimal for cutting manufacturing and is only surpassed by aluminum alloys. Leadcontaining brasses are preferred in precision engineering and the mass fabrication of turning pieces. In addition to the strength, the addition of zinc also affects the color. With increasing zinc content, the color of tombacs with 5-15% Zn moves from golden red (CuZn5) via golden yellow (Cu-Znl5) and yellow (Cu-Zn37) to red yellow when the (3mixed crystal (Cu-Zn40) appears. Some applications for the two differrently processed Cu-Zn alloys: deformable

6.4 Properties and Behavior of Copper Alloys

brass and cast brass are given in the following sections. 6.4.1.1 Deformable Brass Deformable brass alloys have a wide alloying range of about 5-37% Zn, contain the oc-phase, and are supplied in the forms of sheets, rods, profiles, tubes, and wires (Table 6-13): • Cu-Zn5-15 alloys are tombacs with very good cold deformation, very high resistance against atmospheric corrosion, and good conductivity (uses: electrical and metals industries, jewellery, and watches). • Cu-Zn20-33 alloys are also tombacs with good cold deformability, particularly deep drawing (uses: metal tubes, musical instruments, and deep-drawn parts). • Cu-Zn36-37 alloys are the main alloys for cold deformation, consist of the ocphase, are suitable for welding and soldering, and are corrosion resistant (uses: pressure parts, lamps, clamps, zippers, contacts, and sockets). • Cu-Zn40 is good for cold and hot deformation and consists of a + P phase (uses: fittings, hot-pressed parts, and watch housings). • Cu-Zn36-44-Pbl-3 alloys are good for cutting manufacturing, have good hot deformability, and consist of oc + P phase (uses: bottom plates and wheels for watches, parts in precision mechanics and optics, thin extruded parts, and turning pieces). • Cu-Zn20-37 alloys can contain additional alloying elements (1 - 2 % Al, 1 % Sn, 2% Mn, 2% Ni, and others). The alloys are mainly used for construction materials, bearings, chemical engineering, valves, and heat exchangers.

307

6.4.1.2 Cast Brass Cast brass alloys contain 36-43% Zn and 1-3 % Pb: special cast brasses contain, in addition, a few percent of nickel, tin, and manganese. Cast brasses are brittle and cannot be deformed, so the heterogeneous a + P structure is not disturbed; hence they are used when higher strength is required. They have a small melting range and are normally free from segregations. Cast brass alloys can be used for gas and water fittings, furniture fittings, in the electrical industry, and for all castings (Table 6-18): • G-Cu-Znl5 (where G == sand-mold cast) has medium conductivity, good resistance against seawater, and can be soldered (uses: fittings in mechanical, electrical, and precision mechanical industries, and optics). • G-Cu-Zn33-37-Pb alloys have good corrosion resistance, cutting suitability, and conductivity (uses: all pieces and fittings in the sanitary, electrical, and precision industries). • G-Cu-Zn34-37-All-2 alloys have medium strength and cutting suitability (uses: pressure nuts, sliding components, stuffing boxes, ships' propellers, and valve parts). • G-Cu-Znl5-Si4 has high corrosion resistance and can be welded (uses: highly strained, thin parts for the mechanical, naval, and electrical industries). The dissolution of zinc (dezincing) out of Cu-Zn alloys is a form of local corrosion that appears in aqueous solutions. With increasing dissolution of zinc and copper, the electrochemical potential shifts in such a way that the more noble copper is reprecipitated out of solution, while the less noble Cu-Zn-mixed crystal corrodes further. Also, in a + p alloys, the less noble

308

6 Copper-Based Alloys

Table 6-13. Physical and mechanical properties and some applications of Cu-Zn, Cu-Sn, Cu-Al, and Cu-Ni deformable alloysa>b. Alloy

Density Q

Linear Electrical Thermal Tensile thermal conducconduc- strength, expansion, a tivity, tivity, Rm (25 - 300 °C) x (20 °C) k (20 °C) (g/cm3) (10" 6 /K) (m/Qmm 2 ) (W/m K) (N/mm2)

Brinell hardness, HB10

Cu-ZnlO

8.8

18.2

24.7

184

240-350

50-105

Cu-Zn20

8.7

18.8

19.0

142

270-490

55-145

Cu-Zn37

8.4

20.2

15.5

121

290-610

55-190

Cu-Zn37-Pb0.5

8.5

20.4

14.7

113

290-540

60-160

Cu-Zn39-Pb2

8.4

21.1

13.9

109

360-590

85-160

Cu-Zn44-Pb2

8.4

21.2

16.4

126

380-610

90-180

Cu-Zn28-Snl

8.5

19.5

14.1

109

>320

65-100

Cu-Zn35-Ni2

8.3

18.0

5.7

46

440-540

120-150

Cu-Zn37-All

8.3

21.1

7.8

80

440-510

125-145

Cu-Sn6 Cn-Sn8

Q Q

o.o

8.8

18.5 18.5

9.0 7.5

75 67

350-650 370-690

80-200 90-210

Cn-Sn6-Zn6

8.8

18.4

9.5

80

610-760

190-230

Cu-A18

7.7

17.0

8.0

67

370-490

90-130

Cu-A19-Mn2

7.5

17.0

6.5

54

490-590

110-150

Cu-A110-Ni5-Fe4

7.5

17.0

6.0

50

640-740

180-240

Cu-Ni9-Sn2

8.9

17.6

6.4

48

340-560

75-190

Cu-NilO-Fel-Mn

8.9

17.0

5.3

46

420-450

70-120

Cu-Ni25 Cu-Ni30-Mnl-Fe

8.9 8.8

15.5 16.0

3.1 2.7

29 29

^290 ^350

70-100 80-120

Cu-Ni44

8.9

13.5-15

2.0

23

420

85-115

"Average values depending on the heat-treatment and deformation;

b

Applications

tombacs, electrical conductors, jewellery metal tubes, deep drawing clamps, contacts, sockets extruded profiles, wheels, plates precision parts, optical parts thin, extruded profiles ribbed pipes, heat exchangers apparatus, ship construction sliding bearings and elements springs, contacts, pipes chemical apparatus, teeth springs, membranes seawater, valves, apparatus bearings, gears, valve seats condensators, chemical industry spring contacts, plugs, switches seawater, ribbed pipes, plates, apparatus coins, claddings seawater, pipelines, condensors constantan, resistors, wires

According to DKI (1991, 1992).

309

6.4 Properties and Behavior of Copper Alloys

(3-phase will corrode first. These problems can be avoided by the use of special brasses.

60

6.4.1.3 Special Brasses Special brasses is the name given to CuZn alloys with additional alloying elements, e.g., aluminum, iron, manganese, nickel, silicon, and tin in the range 0 . 1 10%. Additional elements are added to improve the corrosion resistance, the strength, and especially the high temperature strength or further characteristic qualities. Aluminum increases the strength and improves the resistance to corrosion and oxidation. Arsenic and phosphorus also improve the corrosion resistance. Phosphorus is also an advantage in cast alloys, as it improves the flow of the melt. Iron and manganese improve the bearing properties and refine the grain size. Nickel improves the high-temperature strength and the corrosion resistance. Tin and especially silicon improve the bearing properties and the corrosion resistance. Special brasses are mainly used for pressure nuts, ships' propellers, valves and control components, bearings, hard solders, and cartridges. The strength of special brasses ranges from 300 to 790 N/mm 2 and the fracture elongation from 45 to 10%. Special brasses can also be cast and soldered. The applications, alloys, and special characteristics are numerous (Wieland, 1986) (see Table 643). 6.4.2 Copper-Nickel-Zinc Alloys (Nickel-Silvers) Nickel-silver alloys are Cu-Zn alloys in which the copper content is partly replaced by nickel, and they contain 7-25% Ni, 11-42% Zn, and 45-62% Cu (see Fig. 6-21) (USN numbers C75 000-C79 999;

30 Ni (wt.%)

50

60

Figure 6-21. Copper corner of the ternary Cu-Ni-Zn system; isothermal section at room temperature with isothermal lines in the liquidus plane (liquid melt). The hatched area indicates common alloys, the straight broken lines are isotherms and the solid curves phase boundaries (DKI, 1991b).

DKI, 1991c; Wieland, 1986). This family of materials has a silver-white color because of the nickel addition, and is therefore called nickel-silver, Neusilber, or alpaca. The Cu-Ni-Zn alloys are characterized by high corrosion resistance, high strength (Fig. 6-22), and good elastic (spring) stability. They are also used at low temperatures, where they retain their ductility, and they are nonmagnetic. As well as additions of manganese (0.4%), iron (0.1-5%), or aluminum (0.5-2%), they occasionally contain lead (1-3 %) for better machining. Brass and nickel-silver are susceptible to stress-corrosion cracking. Components which are under stress or have internal stress are prone to cracking in the presence of moisture, especially when there are traces of ammonia and moisture. This can be avoided by annealing at 200-300 °C (stress-free annealing, soft annealing).

310

6 Copper-Based Alloys

try, in the naval, food, and precision engineering industries, or for coins (Table 6-14). 6A.2.1 Deformable Alloys • Cu-Nil2-Zn24, Cu-Nil8-Zn20, and Cu-Ni25-Znl5 (see Table 6-14) are corrosion resistant and suitable for cold deformation (uses: cutlery, deep-drawn and precision mechanics parts, arts and crafts, relays, contacts, fittings, and membranes). They retain their color. • Cu-NilO-Zn42-Pb, Cu-Nil2-Zn30-Pb, and Cu-Nil8-Znl9-Pb are lead-containing alloys which are good for machining and hot pressing (uses: precision parts, spectacles, zippers, keys, and the optical and watch industries). 0

10

20 30 40 50 60 70 Rolling deformation (%)

80

6.4.2.2 Cast Alloys

Figure 6-22. Influence of cold deformation on the mechanical properties of Cu-Ni25-Znl5 (DKI, 1991b).

Cu-Ni-Zn alloys are used in applied arts, for cutlery, jewellery, musical instruments, or in furniture construction. In addition, they are used for resistors, plugs, and switching springs for the electrical indus-

• ASTM B149 alloys 10A, 11A, and 11B (Table 6-14) are standardized cast alloys of the Cu-Ni-Zn-Pb-Sn family (UNS numbers C97 300-C97 999). They are very good to use for castings, drilling, and machining. They retain their color and are corrosion resistant (uses: armatures, fittings, valves, cast arts, and toilet articles).

Table 6-14. Physical and mechanical properties of Cu-Ni-Zn alloysa Designation

Density (g/cm3)

Linear thermal expansion, a (25- 300 °C) (10" 6 /K)

Cu-Nil2-Zn24a Cu-Nil8-Zn20 Cu-Ni25-Znl5

8.7 8.7 8.8

16.4 17.0 16.6

4.0 3.4 2.8

Cu-NilO-Zn42-Pb Cu-Nil2-Zn30-Pb Cu-Nil8-Znl9-Pb

8.5 8.6 8.8

19.4 19.6 16.9

5.3 4.2 3.4

10A, cast b 11 A, cast 11B, cast

8.9 8.8 8.8

Q

a

Deformable alloys DIN 17663;

Electrical Thermal conductivity conductivity at 20 °C, x at20°C, A (m/Q mm2) (W/m K)

(N/mm2)

Brinell hardness, HB 2.5/62.5

33 27 23

340-610 370-610 390-540

85-185 90-185 90-165

35 33 27

510-590 490-590 430-530

150-170 155-175 135-160

210 210 310

50- 60 90 135-150

3.4 3.0 3.8 b

Tensile strength,

Cast alloys ASTM B149, average values.

6.4 Properties and Behavior of Copper Alloys

311

6.4.3 Copper-Tin Alloys (Bronzes) Cu-Sn alloys are the classical bronzes. They belong to the oldest class of alloys used by mankind, and have great importance in the history of civilization. They were already known in the so-called Bronze Age, and since about 2700 B.C. the tin bronzes have been used for weapons, vessels, utensils, jewellery, and coins. Alloys for mechanical forming are limited to a maximum of 9 % Sn (UNS numbers C50100-C54 899; DIN 17 662), while cast alloys normally contain between 10 and 20% Sn (DIN, 1991; UNS numbers C90200-C94999; and DIN 1705). With an increase of the tin content, the wear and chemical resistance increase. Even under high strains, bearings show good tribological properties. According to the binary phase diagram (Fig. 6-23; Hansen and Anderko, 1958; DKI, 1992b), the tin bronzes with up to

liquid melt /3 + liquid melt

°.h

Figure 6-24. Microstructure of continuous casting GC-Cu-Sn5-Zn-Pb for bearings; a-phase with Pb inclusions.

14% Sn solidify as the ductile a-phase (f.c.c.) (Fig. 6-24), but substantial segregations appear as a result of the very wide solidification range. The p- and the 8phases are brittle, hard, and difficult to deform. Also, diffusion is very slow below 520 °C. Therefore Cu-Sn alloys with more than 6 % Sn consist of the oc-phase and the eutectic a + 8 (through technical cooling) after solidification, which results in a strong reduction in their hot deformability because of the brittle 8-phase. Before casting, tin bronzes must be deoxidized with phosphorus, and as such a brittle, phosphorus-containing phase frequently appears in the microstructure (Bargel and Schulze, 1994). Phosphorus enhances the brittleness and increases the melt flowability and the corrosion resistance. 6.4.3.1 Deformable Alloys

10

20 Sn (wt%)

Figure 6-23. Copper corner of the Cu-Sn phase diagram. The solid lines correspond to normal annealing time (homogenizing) and the broken lines to technical cooling (fast) (DKI, 1992 b).

Deformed Cu-Sn alloys have tensile strength 350-700 N/mm 2 , yield strength 290-600 N/mm 2 , hardness (HB) 90-210, elongation (A5) 20-7% (Fig. 6-25); electrical conductivity at 20 °C 7.5-12 m/Q mm 2 (Fig. 6-26), and thermal expansion (20-200°C) 18.2-18.5 x 106 1/K (DKI, 1992 b). They are used for springs, con-

312

6 Copper-Based Alloys

I 800 r
200 2

o

4 6 Sn (wt.%)

I 800

250

4 6 Sn (wt.%)

r

CD

z <M

2

Cu

600

u. 200 -

400

"2 150

ai ro

200

0 Cu

Undeformed

Undeformed

=3 100

4 6 Sn (wt.%)

2

4 6 Sn (wt.%)

Figure 6-25. Mechanical properties of Cu-Sn alloys with deformation in relation to the tin content (DKI, 1992 b).

tacts, and conductive strips (Table 6-13): Cu-Sn2 and Cu-Sn5 are used for screws, conductive springs, contacts, plugs, wire wraps, and thermi-points. Cu-Sn4-Zn4 and Cu-Sn4-Zn4-Pb4 are used for conductive springs, strips, and bands.

20°C,0.05%P 200°Cl0.05%P 20°C,0.40%P

4 6 Sn (wt.%)

Figure 6-26. Electrical conductivity, x, of Cu-Sn alloys between 20 and 200 °C in relation to the tin and phosphorus content (DKI, 1992 b).

6.4.3.2 Cast Alloys Cast Cu-Sn alloys with 10-12% Sn are used for worm gears, gear wheels, slide bearings, high-pressure fittings, and pump wheels. They have the following properties: Rm = 250-300 N/mm 2 , Up0>2 = 110150N/mm 2 , v45 = 30-18%, and Brinell hardness, HB = 70-100 (Weisner e t a l , 1982). They have the following applications: • G-Cu-Snl2, G-Cu-Snl2-Ni, and G-CuSnl2-Pb are used for construction. • G-Cu-Snl4 with its maximum strength and sufficient toughness is used for gear wheels and highly-strained construction parts. • G-Cu-Sn20 and G-Cu-Sn25 are hard and brittle because of the oc + 8 eutectic, and are used for bell castings. For cast Cu-Sn alloys, increased strength of the a-mixed crystals (f.c.c.) is obtained

313

6.4 Properties and Behavior of Copper Alloys

by further additions of tin of up to 14%, while the microstructure has a large proportion of the brittle 8-phase (b.c.c.) caused by technical (quick) cooling, which must be removed by heat-treatment. 6.4.3.3 Red Castings

Red castings are Cu-Sn alloys with additional zinc and lead. They contain 3-11 % Sn and 2-15 % Zn, both of which improve the casting characteristics, reduce the possibility of oxygen absorption by the melt, and reduce the cost. However, zinc reduces the strength and the deformability. A content of 1-5% Pb enhances the sliding behavior and the machinability. An addition of 1 % Ni leads to smaller grain sizes and improved toughness. Red castings are used for paper rollers, slide bearings, fittings, ships, shaft claddings, or sculptures. Their properties include: Rm = 280-270 N/mm 2 , A5 = 1518%, HB = 60-75, and x = 7.5m/Q mm 2 (Weisner et al., 1982). Majors specific uses: • G-Cu-SnlO-Zn, G-Cu-Sn7-Zn-Pb, and G-Cu-Sn5-Zn-Pb (Fig. 6-24) are used as sliding materials for bearings. • G-Cu-Snl0-Zn20 is used for sand casting of corrosion-strained, mechanical parts. • G-Cu-Sn8-Zn7-Pb3 and G-Cu-Sn4Zn4-Pb4 are used for bearing and roller shells.

which increases their resistance against seawater, sulfuric acid, and salt solutions, as well as giving them good wear properties and oxidation resistance. The excellent corrosion resistance of Cu-Al alloys is based on a protective layer of alumina, which builds up quickly and can heal when damaged (DKI, 1991b). The material with more than 10% Al cannot be deformed. According to the phase diagram (Fig. 6-27), the homogeneous oe-phase (f.c.c.) exists up to 9% Al (Fig. 6-28). The P-phase can be transformed at temperatures below 565 °C into eutectic (oc + y) structure (equilibrium) or into a martensitic structure by quenching: Further alloying elements are iron (grainrefining, magnetic permeability), manganese (deoxidation), nickel (excellent corrosion resistance), arsenic (resistance against salt solution), and silicon (elevated strength, better cutting). Because of their good (high-temperature) strength in the hot deformed and cast condition (Fig. 6-29), Cu-Al alloys are

Liquid melf

6.4.4 Copper-Aluminum Alloys

Aluminum bronzes are the copper alloys with the best chemical resistance; they also have high strength. They are standardized as deformable material in UNS numbers C60 600-C64499 and DIN 17 665, and as cast material in UNS numbers C95 200C95 900 and DIN 1714 (CDA, 1982; DKI, 1991b). Cu-Al alloys contain 3-14% Al,

12

14

Figure 6-27. Copper corner of the binary Cu-Al phase diagram (Hansen and Anderko, 1958; DKI, 1991b).

314

6 Copper-Based Alloys

Figure 6-28. Microstructure of continuous casting of GC-Cu-A19.5-Ni; a-structure with eutectic at grain boundaries.

s80 ^ |

70 60

I «> £

30

I 20

Cu-A15-As is resistant against salt solutions (use: heat exchangers). Cu-A18 and Cu-A17-Si2 are resistant against seawater (uses: valves, screws, and bearings). Cu-A18-Fe3, Cu-A110-Fe3-Mn2, and Cu-A19-Mn2 have good high-temperature strength, are resistant against erosion and cavitation, and are shockproofed against stress (uses: heat exchangers, nuts, gears, bearing shells, and valve seats). Cu-A19-Ni3-Fe2, Cu-A110-Ni5-Fe4, and Cu-Alll-Ni6-Fe5 are resistant against corrosion, friction, and cavitation. G-Cu-AUO-Fe, G-Cu-A19-Ni and GCu-AUO-Ni are good casting and construction materials with high strength and resistance against corrosion and wear.

Sand casting Mold casting Hot rolled

6.4.5 Copper-Nickel Alloys ^

600

~

500

^

400

If 300

"S 200 -^ 100 g

0

0 1 2 3 4 5 6 7 8 9 10 11 12 Al (wt.%) Cu Figure 6-29. Tensile strength, Rm, and fracture elongation, A 5 , of Cu-Al alloys in cast and rolled conditions (DKI, 1991b).

used in steam fittings, guides, and valves. Their increased corrosion resistance enables their application in the food, chemical, and petroindustries, in ship construction, desalting plants, off-shore technology, in the electrical industry, power plants, and for coins (Table 6-13):

The Cu-Ni alloys (nickel bronzes) usually contain between 2 and 45 % Ni, and, in order to improve their special properties, some manganese, iron, and tin, and for cast alloys additions of niobium and silicon. They are standardized as deformable alloys (DIN 17664 and UNS numbers C70100-C71999), as cast alloys (DIN 17658 and UNS numbers C96200C96999), and as resistor metals (DIN 17471; DIN, 1991; CIDEC, 1972). Copper and nickel are very close in the periodic table and have some similar properties (lattice spacing, f.c.c. structure). Therefore they form a complete, homogeneous, oc-mixed crystal (Fig. 6-30) over the complete range of the phase diagram (Fig. 6-31), if alloyed together. As a result of this, many of the physical properties change continually with concentration of

6.4 Properties and Behavior of Copper Alloys

Figure 6-30. Microstructure of forged flange of CuNilO-Fel-Mn; a-phase with coarse and fiber-like. elongated grains.

1600 1500

Liquid melt

1453°C

1400 1300 1200 ?1100 1083°C fcJiOOO

400

358°

300

Paramagnetic

200 100 -

''Ferromagnetic

0

i

0

Cu

10 20

30 40

50

60

Ni (wt.%)

70 80

i

1

90 100

Ni

Figure 6-31. Binary Cu-Ni phase diagram and magnetic condition (Hansen and Anderko, 1958; DKI, 1992 a).

alloying element, this is the basis of many applications in electrical engineering. Even a small amount of nickel (1.5% Ni) doubles the impact toughness of Cu-Ni alloys compared to copper, and at the same time the corrosion resistance is improved. The high-temperature strength of the corrosion-resistant Cu-Ni alloys is comparable to that of stainless steel (DKI,

315

1992 a). The color of the copper alloy changes with increasing nickel content; from about 15% Ni the Cu-Ni alloys are. nearly silver-white. The physical and mechanical properties of Cu-Ni alloys are influenced by many additional alloying elements. Manganese is used for deoxidation of the melt and improves the casting properties, strength, warm-hardening, and temperature of strength degradation. Iron dissolves in the a-phase and can increase the corrosion resistance by forming a strong, passive layer (1.5% Fe and 2% Mn). Tin addition improves, above all, the strength (Cu-Ni9Sn2), resistance to color change, and wear properties, and has a good spring stability. Silicon gives better casting flow, high-temperature strength, and ductility. For alloys with a low nickel content (0.5-4% Ni), further alloying with silicon (0.2-1% Si) results in good age-hardening due to the formation of Ni2Si. The lead content must be low (under 0.02% Pb) in deformable alloys and for welding, but it is used in cast alloys to aid cutting (1-11 % Pb). Zinc is a main constituent of nickel-silver (see Sect. 6.4.2). Aluminum, beryllium, and chromium are used for high strength and seawater applications. Cu-Ni alloys have two main fields of application: they are used because of their special electrical properties and as corrosion-resistant materials. 6.4.5.1 Special Electrical Applications Special electrical applications of Cu-Ni alloys are as wires, sheets, strips, profiles, and layers of resistor materials, e.g., CuNi44 (Konstantan). The electrical resistivity of copper is increased by additions of alloying elements, even nickel. This reaches a maximum at about Cu-Ni50 (Fig. 6-32). The strength and hardness also change in

316

6 Copper-Based Alloys

0 Cu

20

40 60 Ni (wt.%)

80

100 Ni

Figure 6-32. Electrical resistivity, @e, and coefficient of thermal expansion, a, for Cu-Ni alloys at room temperature (resistor metals) (Guillery et al., 1991).

the same way across the composition range of the Cu-Ni system; the strength increases for both elements reaching a maximum near the middle of the system. The thermo-voltage against copper or iron also reaches a maximum at 45 % Ni, therefore Cu-Ni44 is used as a material for thermocouples (Cu/Cu-Ni44 or Fe/CuNi44). The resistor materials are divided into three groups: resistors for precision measurements (Qe(T) constant), resistors for controlling (QQ(T) variable), and resistors for heating and energy (ge high). • Resistors for precision measurements have an electrical resistance Qe« 0.4 ^Q m, which is constant with temperature; they consist of copper alloyed with nickel and manganese, e.g., CuMnl2-Ni2 (manganin), Cu-Mnl2-A14Fel.5 (novoconstant), and Cu-Mn-Sn. • Resistors for controlling also have QG « 0.4 JIQ m at room temperature, but this alters with temperature. They are used as starting resistors in vehicles and lamps and as rheostat resistors for igni-

tion; they include Cu-Ni-Mn alloys, Cu-Ni20-Zn20 (nickel-silver), CuNi45-Mnl (constantan), Cu-Ni30-Mn3 (nickeline), and Cu-Mnl3-A13 (isabellin). • The third group, heating and energy resistors, must have a high resistance Qe « 0.9-1.5 JIQ m in order to produce heat, and are used in furnaces, household utensils, and brakes; they consist of iron alloyed with chromium, nickel, and aluminum, e.g. Fe-Cr30-Al, Fe-Ni30. 6.4.5.2 Corrosion-Resistant Alloys Corrosion-resistant Cu-Ni alloys are construction materials of different semiproducts. The corrosion resistance of copper can be improved considerably by alloying with nickel, which at the same time increases the strength. This is the result of passive layers, which are formed in the presence of sufficient oxygen. These Cu-Ni alloys, normally used with between 5 and 30% Ni (Fig. 6-33) and alloyed with iron or manganese, are resistant to seawater and are used with other alloys in condensers, in seawater desalting plants (sludge fittings), and on ships. However,

6oo r

10 15 20 Ni (wt.%)

Figure 6-33. Tensile strength, Rm, and fracture elongation, A5, of Cu-Ni alloys (DKI, 1992a).

6.4 Properties and Behavior of Copper Alloys

there are some welding problems with these alloys resulting in porous welds. • Cu-Ni25 is used for many coins. 6.4.5.3 Deformable Alloys

and cavitation, and are good for welding and cutting (uses: paper machinery, ship construction, chemical engineering, pumps, fittings, stirring and power plants, and the food industry).

Deformable alloys have the following properties and applications (Table 6-13):

6.4.6 Copper Alloys with Silicon, Beryllium, Manganese, or Lead

• Cu-Ni9-Sn2 has good cold deformability and is spring hard (uses: spring contacts, relays, plugs, and lead frames for semiconductors). • Cu-NilO-Fel-Mn is very resistant to erosion, cavitation, and corrosion (uses: seawater application of pipes, sheets, condensers, heat exchangers, brake linings, offshore structures, and rib pipes). • Cu-Ni25 is wear resistant (uses: coins and claddings). • Cu-Ni30-Mnl-Fe and Cu-Ni30-Fe2Mn2 have improved strength and corrosion resistance (uses: pipelines and condensers in seawater). • Cu-Ni44 and Cu-Ni44-Mnl (uses as for Cu-Ni30-Mnl-Fe).

6.4.6.1 Silicon Bronze

6.4.5.4 Resistor Metals • Cu-Ni2, Cu-Ni6, and Cu-NilO have low electrical resistance and can be soft soldered (uses: precision resistors and heating wires). • Cu-Ni23-Mn and Cu-Ni30-Mn are corrosion and scale resistant (uses: electrical resistors, starters, and rheostat resistors). • Cu-Ni44 and Cu-Ni44-Mnl have good hot and cold deformability (uses: resistors for controlling, thermo-elements, and heating wires). 6.4.5.5 Cast Alloys • G-Cu-NilO and G-Cu-Ni30 have good resistance against corrosion, erosion,

317

Two particular Cu-Si alloys are important, i.e., Cu-Si3-Mnl and Cu-Si2-Mn0.5. Silicon has a deoxidizing effect and increases the strength and corrosion resistance; manganese improves the high temperature strength and wear resistance (Schimpke etal., 1977). These materials are used particularly in the refrigeration industry, in heat exchangers, armatures, and chemically resistant castings. 6.4.6.2 Beryllium Bronze Beryllium bronze contains 0.6-2.8% Be. The material, which is easy to form cold, can be hardened by the precipitation of Cu2Be. The phase diagram (Fig. 6-34) shows three different phases for the normally used Cu-Be alloys: oc-phase (f.c.c, good deformability), (3-phase (b.c.c, eutectic decomposition), and y-phase (b.c.c, intermetallic, and brittle). In copper alloys beryllium increases the mechanical properties and the hardenability, but decreases the electrical and thermal conductivities. The beryllium content can be reduced by replacement with alloying elements such as zirconium, cobalt, silicon, or silver. Normally some nickel (0.5% Ni) or cobalt is added to reduce grain growth, silver to prevent surface oxidation and to increase the conductivity, and lead to improve the machinability (Heller, 1976). Cu-Be alloys are the hardest of all the copper alloys. They have high strength and

318 Cu 1100

10

6 Copper-Based Alloys

20

30

Be (at.%) 40 50

in the electrical industry, e.g., overhead wire for railways, antennas, relay contacts, leaf springs, electrodes for soldering machines, roller bearings, gears, die molds, antimagnetic materials, and spark-proof tools.

60

• Cu-Bel.7 and Cu-Be2 have high strength, wear resistance, and good conductivity (Table 6-15). • Cu-Co2.5-Be0.5 and Cu-Col.5-AglBe0.4 have high conductivity and good strength.

8

10 12 Be (wt.%)

14

16

18

20

Figure 6-34. Copper corner of the binary Cu-Be phase diagram (Hansen and Anderko, 1958; Dies, 1967).

conductivity (Fig. 6-35). The most important and commonly used alloy is Cu-Be2 with the highest strength, i?m = 1500N/ mm 2 in the hardened condition (about three times that of other copper alloys). For applications requiring good electrical conductivity, Cu-Co-Be (2.4-2.7% Co, 0.4-0.7% Be) is suitable with its conductivity, x = 28 m/Q mm 2 (about half that of copper). As the electrical conductivity in the hardened condition is still 10-20 m/fi mm 2 , these Cu-Be alloys are mainly used 50 r -

6.4.6.3 Manganese Bronze Manganese bronzes are resistant to seawater and are very stable. Cu-Mn alloys are used as casting materials and for ship propellers. 6.4.6.4 Lead Bronze Lead bronzes contain 5-28% Pb, 4 10% Sri, and small additions of nickel, zinc, iron, and antimony (Schimpke et al., 1977). They are standardized in DIN 1716, and are good for metal cutting, drilling, etc. Of all slide bearing materials, they have the highest heat transfer ability as well as good emergency performance. • Cu-Pb25 is used for castings and especially for bearings with dynamic stress (diesel engines).

1500 r

Figure 6-35. Influence of annealing temperature on the tensile strength, Rm, modulus of elasticity, E, elongation, 8, electrical conductivity, x, and Brinell hardness, HB, of cold-rolled Cu-Be2; annealing time: 10 h (100% IACS = 58.0MS/m). 300 Temperature

319

6.5 Technology and Products of Copper-Based Alloys

Table 6-15. Physical data of copper-beryllium alloys (in the hardest condition after precipitation hardening) compared with copper (oxygen-free high-conductivity)3. Property

Symbol

Melting point (solidus) Density Modulus of elasticity Shear modulus

T

Electrical conductivity at 20 °C

X

Electrical resistivity at 20 °C Thermal conductivity at 20 °C Linear expansion coefficient at 20-200 °C Specific heat capacity at 20 °C Yield strength Tensile strength Elongation Hardness (Vickers) Fatigue strength for 108 cycles 1

Q

E G

X a C

P

^p0.2

3 HV

OFHC-Cu

Cu-Be2

Cu-Co2.5Be0.5

Cu-Col.5Agl-Be0.4

°C g/cm3 N/mm 2 N/mm 2

1083 8.96 125 000 46000

865 8.25 130 000 51000

1030 8.75 133 000 52000

8.72 132000

MS/m % IACS fxQm W/mK

59 101.7 0.017 393-460

12.7 22 0.078 113-126

27.5 47.5 0.038 209-239

28.5 49.1

1/K J/gK

16.2 xlO~ 6 0.385

17.0 xlO" 6 0.418

17.6 xlO" 6 0.418

17.6 xlO" 6

N/mm 2 N/mm 2 % N/mm 2 N/mm 2

40- 430 150- 450 5 - 40 400-1000 <180

200-1200 450-1500 5 - 35 900-4000 250- 340

140- 900 390-1000 10- 25 700-2600 210- 250

140- 840 250- 910 8 - 25 750-2600

Unit

Heller (1976).

6.5 Technology and Products of Copper-Based Alloys Copper is the main material used in electrotechnology because of its combination of material properties, especially its excellent electrical conductivity, in addition to sufficient mechanical strength and good corrosion resistance. Today more than half of the total copper production is used in the electrical industry (Fig. 6-5). This tendency is growing slightly (Metallgesellschaft, 1993; DKI, 1982). In relation to special requirements and applications, the final and semiproducts of copper alloys are divided into five groups: (1) Highly conducting copper alloys, which retain the good conductivity of copper as much as possible, while increasing the mechanical strength. (2) Construction materials with high corrosion resistance and essential improved

mechanical strength, as a result of which a large decrease of the conductivity occurs. (3) Contact materials with low contact resistance, due to their good conductivity and corrosion resistance. (4) Resistance materials with intentionally high specific electrical resistance, which is as independent of temperature and environmental influences as possible. (5) Other electrical applications, which are very specialized. 6-5.1 Applications in Electrical Engineering and Electronics

For copper alloys and components, which are applied in the electrical and electronic industries, high conductivity values, a favorable bending and softening behavior, as well as good workability are required more and more. All these conditions cannot, however, be fulfilled by any one material. For this reason, a number of

320

6 Copper-Based Alloys

different copper alloys have been developed, whose properties meet the respective requirements and hence satisfy various application ranges. Copper is initially manufactured in the unalloyed condition as cables, wires, ropes, rails, bands, strips, and other components for the electrical industry. As a result of the increased miniaturization of electronics, the growth of copper uses in this field is not so high; but in the field of electrical energy further increases are expected in the future (Arpaci and Bode, 1992). Some of the typical and numerous applications of copper and copper alloys are summarized below for the main fields of electrotechnology. 6.5.1.1 Electrical Conductors

For electrical conductors the combined properties are of interest; besides the very high conductivity, the advantages of high wear and corrosion resistance, low magnetic permeability, the fact that sparks are not struck, adequate strength, and good heat transfer in many different forms are used. As a conducting material copper also has an optimal price to value relation (Metallgesellschaft, 1993). Cables for installation and conduction are mostly made of E-Cu58 and E-Cu57, but also of OF-Cu, SE-Cu, Cu-AgO.l, and Cu-AgO.l-P. Here two different applications are preferred, first as massive copper wires for locally fixed conductors, second as bundles of fine strands for movable uses. The cables with fine strands are suitable for uses subject to recurrent high bending stresses (e.g., conduction for elevators, conduction for coils, and special elastic insulation cables). Cables and conductors for high current and used for energy transmission and dis-

tribution (e.g., community distribution, large industries, supplying large consumers, dredgers, mining machinery, cranes, and welding machines) are mainly produced from C-Eu58 and E-Cu57. These cables for fixed installation are used with wire of cross section up to 16 mm 2 max. (4.5-mm diameter), while fine strand wires are used with-cross section of up to about 630 mm 2 (28 mm diameter). For the sake of good contact, they are coated with a thin layer of tin or silver (no oxide formation at the surfaces). Conduction ropes and rail wires are made from conductive bronzes because of their high strength (Table 6-11), and include, e.g., Cu-Mg0.4, Cu-AgO.l, Cu-Ni2Si, and E-Cu58 (Table 6-10); they also have different profiles for good fixation. For electrical rail wires massive, drawn wires or extruded profiles (round, groove) are used. In conveyor equipment round driving wires of E-Cu57 are mostly applied. For earthing systems (e.g., protective conductors), copper rods and strips up to a cross section of 50 mm 2 are mostly used, or copper-cladded bands of steel with 20 % copper cladding, or steel with a zinc layer. A high corrosion resistance is achieved for earthing conductors by using steel with a copper coating, copper with tin and zinc coatings, or pure copper. Coils for large machinery, transformers, generators, and motors are fabricated from drawn rectangular or hollow profiles (the inside is used for cooling by water or gas) and of E-Cu57 in a soft condition. Conductors for frequent, substantial alteration of the temperature or for high temperatures are preferentially made of lowalloyed copper with Ag, Cr, or Zr (e.g., Cu-Ag0.03) because of their good high temperature strength (Table 6-10) (Guillery etal., 1991).

6.5 Technology and Products of Copper-Based Alloys

Coils are produced from round wires for various pieces of apparatus and equipment (e.g., stator and rotor coils of motors, electromagnets, relays, choke coils transformers for current, voltage, and energy transmission). These wires are insulated by glass, paper, or lac and have a diameter from 0.02 to about 6 mm in the round, flat, or rectangular profile of the conductor. E-Cu58 (Table 6-10) is used in the soft condition for these applications. Communication network conductors (low and high frequency) are made of symmetric wire pairs (two wires), normally with 0.4-0.8 mm diameter. All these conductors and cables are usually made of ECu58. For telecommunications, copper alloys are the preferred conductors (conductive bronzes) besides optical fiber glass. For the transmission of high frequencies, coaxial wires or hollow conductors are necessary in many cases; microwaves are conducted in rectangular hollow profiles of Cu-Zn30 or in tubes of E-Cu57 or E-Cu58. The conducting network of integrated circuits (IC) is made of pure copper with a thickness of 10-100 jim. To improve the integration, conducting plates are mounted on both sides (surface-mounted devices, SMD techniques), a multilayer coating with a metal core (e.g., Cu-Invar-Cu) is built up (Arpaci and Bode, 1992), or copper metallization is used (see Sec. 6.6.7). The electrical conducting contact can be made by mechanical fixation or by plugs, or by soldering and welding copper and copper alloys (Steffens and Sievers, 1990; Ruge, 1991). Because of heat transfer and thermal expansion, the welding and soldering of copper is difficult and can produce impurities and brittleness inside the material.

321

6.5.1.2 Construction Materials Construction parts of electrical machines or switch gears are mostly fabricated from copper alloys with a compromise between the conductivity and the strength (e.g., contact carriers, switches, clamp parts, and connecting screws). For lowstrength applications, copper wrought alloys and castings (sand or mold) are used, while for high-strength applications cast Cu-Cr alloys in a hardened condition or forged parts of Cu-Cr-Zr, Cu-Ni2-Si, or Cu-Ni3-Si in the hardened condition are preferred. Conducting lines and rails of switch gears, furnaces, and transformers are produced as round or rectangular wires of E-Cu57, SE-Cu or the alloys Cu-Ag 0.03 or Cu-Ag0.1. Aluminum conductors are mostly coated with copper to achieve an oxidation-resistant surface and a better distribution of the current over the cross section. 6.5.1.3 Contact Materials Contacts are necessary in all circuits, and their electrical resistance must be as low as possible. This can be obtained either by using a surface that is free of oxides or by having a high contact pressure. Static contacts are screwed, welded, or soldered (uses: clamped connections for conducting lines, wires, rails, etc.). Plug contacts (e.g., used in household, automobile, and electronic applications) are made of, e.g., Cu-Zn37, Cu-Zn39-Pb2, or CuSn6, and contact springs are made of CuNi9-Sn2 or Cu-Co2.5-Be0.5 (with high resistance against relaxation). For sliding contacts (e.g., for commutator, motor, or household equipment), unalloyed copper E-Cu57 is used in a hardened condition; alternatively, low-alloyed copper is used for higher thermal and wear

322

6 Copper-Based Alloys

resistance, e.g., the alloys Cu-Ag0.03, CuAgO.l, Cu-AgO.3-P, Cu-Cr, or Cu-Cr-Zr (Guillery et al., 1991), or copper sintered with graphite. Pure copper plug contacts are mainly used in energy techniques with high currents and in different mediums (e.g., air, compressed air, SF 6 , oil, vacuum, etc.). For high-energy contacts some copper compounds can be used, e.g., sintered CuW (with high strength against decomposition) or precipitation-hardened copper alloys such as Cu-Cr. Pure or silvered copper, or Cu-Ag or Cu-Cr alloys are used for separation switches, and for relays and connecting springs Cu-Sn6 or Cu-Be2 are used (with high relaxation and corrosion resistance). Pure copper (E-Cu58 or SE-Cu) is frequently selected as the basic material for contacts with cladded noble metals because of its high conductivity. Cu-ZnlO37 or Cu-Sn6-8 alloys are used in low-current techniques because of their reasonable conductivity and high fatigue strength. 6.5.1.4 Resistor Metals Resistance materials are necessary to produce heat for radiators, furnaces, etc. The wires and coils are fabricated from different copper alloys, e.g., Cu-Ni2, CuNi6, Cu-NilO, Cu-Ni23, and Cu-Ni30Mn, and have an electrical resistivity of about 0.05-0.5 Q mm 2 /m. Resistor metals, e.g., Cu-Ni44 (constantan), Cu-Mnl2-Ni (manganin), Cu-Mnl3A13 (isabellin), and Cu-Mnl2-A14-Fel (novoconstant), are used for measuring the resistance or in automatization for control circuits, up to 400 °C (Fischer, 1978). They have a higher resistivity (above 0.5 Q mm 2 / m) than heating materials, a smaller thermal expansion, and good temperature constancy over a wide range of temperature

with a linear relation between resistance and temperature. 6.5.1.5 Other Electrical Applications Some special applications of copper and copper alloys are summarized here, which depend on particular properties of the material. Semiconductor supports (chip carriers) are made of low-alloyed copper, e.g., CuFe2-P, Cu-Zn0.5, and Cu-Sn0.15, because of their high electrical and thermal conductivity. They contact and conduct the electrical current and act as heat sinks. Superconductors can easily be fabricated as thin filament wires while coated with copper carrier. For mechanical, electrical, and thermal stabilization, high purity copper and Cu-Sn alloys are used around superconducting filament wires. The copper coating transfers the current transportation for a short time by local breakdown to these compound superconductors with up to 6000 internal thin Nb 3 Sn filament wires (multifilament superconductor, drawn to about 5 [am diameter) or to alloyed superconductors with up to 100 Nb-Ti filaments. For stabilization against eddy currents inside the combined superconductor (total diameter about 0.4-1.2 mm), it is possible to use Cu/Cu-Ni carriers around the superconducting filament wires. These filled copper wires can easily be deformed into very thin filaments. Superconductors of these types are used in magnetic coils with high and very homogeneous magnetic fields of up to 10 T. The electrodes for electrical spot welding and welding with a rolling seam of metals must have high electrical conductivity, good heat transfer, and a low welding tendency (no paste over) with the welding material. In addition, high strength and

6.5 Technology and Products of Copper-Based Alloys

hardness are required in order that the high pressing force can be reached, even after a large number of spots, without deformation of the electrode. Therefore copper alloys are used, e.g., Cu-Ni2-Be, CuCr, Cu-Co2.5-Be0.5, Cu-Ni2-Si, Cu-Be2, and Cu-Bel .7, and for lower requirements, E-Cu. In thermoelements the thermo-voltage at the junction of two wires made of different metals or alloys is used for measuring temperature differences, e.g., a thermoelectric couple of E-Cu58 and Cu-Ni44 from low temperatures up to 400 °C, because of the thermo-electric potential of 4.25 mV per 100 °C. Thermal bimetals are used for example in the switches for current circuits, and are released by a change of the temperature. By this method, the two metal strips are differentially expanded by the current flow, so heating the metals and leading to the switch opening. A combination of two metal strips, e.g., Mn-CulO-18-NilO-16 and pure copper, is effective for this. The high deformability of copper and copper alloys makes it possible to deform them into boxes and parts by deep drawing. The protection of mechanically sensitive parts or protection against electromagnetic fields can be provided using boxes or covers that are deep drawn and made from copper alloys, e.g., Cu-Ni9Sn2, Cu-Ni30-Mnl-Fe, Cu-Zn0.5, and mostly Cu-Zn37. These alloys are used for boxes and covers for oscillating quartz, frequency filters, and relays, also for light suspenders, switching contacts, plug contacts, and high-frequency hollow conductors (Arpaci and Bode, 1992). 6.5.2 Copper Semiproducts

In the copper mine or refinery, the metal for final products (e.g., rolled plates, rods,

323

round bolts) is cast as a semiproduct ready for subsequent processing. Nowadays, these cast formats are mainly produced in vertical, continuous or half-continuous extruders, giving semiproducts like sheets, strips, tubes, and profiles. Castings for wires are wire bars, which are produced in open and horizontal molds connected to the air. A new production method is the fully continuous casting-rolling process, by which the cast wire is rolled in one melt. 6.5.2.1 Plates and Strips Plates for rolling are preheated up to about 800-950 °C and hot rolled to a thickness of about 10-20 mm. After quenching, the edges are machined and by further rolling in the cold condition the desired thickness is reached after several rolling steps. In a further step to obtain determined rolling grades, an intermediate soft annealing (soaking) is often necessary to avoid the formation of a strong texture (anisotropy) and accompanying high strength. After heat-treating (soft annealing or hardening) the cold-rolled or colddeformed plate or strip, the material is pickled to obtain a clean and smooth surface. A specified hardness, strength, and quality of surface are required from strips and plates. Finally, the semiproducts are cut to a specific standard length and size, and their material properties are determined by testing a sample. 6.5.2.2 Tubes Copper tubes are produced by different fabrication processes and in various forms. Seamless tubes are made from cast round blocks (clubs), which are hollowed out at a high temperature using extruders with an internal mandrel and pressed as tubes. For the fabrication of pre-pressed tubes another process with diagonal rollers (piercing)

324

6 Copper-Based Alloys

is also used, by which the cast club is heated, deformed, and drawn over the head of the mandrel. This is followed by cold deformation in rolling mills (Jung and Matucha, 1993), raw-drawing equipment, or tube drawing machines (Richter, 1993). An intermediate soft annealing (soaking) is necessary for tubes of alloyed copper and has a remarkable effect on the deformation, surface quality, and size tolerance. Tubes of unalloyed copper can so be deformed by up to 99 % because of its good ductility. Tubes with a noncircular cross section are also made by extrusion with subsequent rolling or drawing. Tubes with longitudinal or spiral welds are made from rolled bands of copper. For the most popular commercial shapes and sizes lead-containing brass is used, e.g., Cu-Zn39-Pb3 (Diehl, 1991), which has good machining properties, is corrosion resistant, and is suitable for hot working and forging. Brass tubes and rods (hot or cold formed) contain additional alloying elements and, as well as a uniform structure, have good machinability, a medium or high tensile strength, improved corrosion resistance, and exceptional tribological properties. Therefore they (and parts made from them) have a large range of applications e.g., bearings, bushes, gear parts, and valve guides. The most common lead-free brass for the production of thin-walled tubes is CuZn37, which can be easily polished, is corrosion resistant, suitable for plating, and uncomplicated for good machining. Therefore these tubes are used for jewellery and furniture products, for plumbing fittings, lighting fixtures and lamps, and also for sanitary installations. 6.5.2.3 Rods and Profiles As in the case of tubes, rods and profiles are fabricated from preheated round bars

or clubs at hot pressing temperatures by hot and cold deformation. Thicker cross sections of rods and extruded profiles are pressed or drawn to their final straight length on drawing machines, or by rolling. Here the lead-containing or lead-free brasses Cu-Zn39-Pb3 or Cu-Zn37 are mostly used. 6.5.2.4 Wires For electrotechnological applications, wires are the most important semiproducts of copper and copper alloys. Cast wire bars or extruded rods are heated to hot rolling temperatures and then reduced in several consecutive rolling steps to an intermediate size (mostly 6-8 mm diameter). These hot-rolled wires are then (after pickling and rinsing) drawn on very fast running, multiple-drawing machines (final velocity up to 1500 m/min) down to a diameter of about 2 mm. To produce very thin wires and wires with a high surface quality, the surface of the pre-drawn wire must be removed before the final drawing operation. A very fine wire can be drawn of down to several one hundreths of a millimeter in diameter (DKI, 1982). Some rolling processes have been economically and technologically improved. The connection of wire ends by cold pressure welding is the basis of continuous fabrication. Reductions in the cross section of up to 99 % are possible by drawing copper wires without intermediate annealing if SE-Cu is used. For very thin wires, intermediate annealing is necessary because of the risk of fracture during drawing. At large copper mines an integrated cast-rolling process of wires is often used. In this case additional heating of the wire bars is not necessary, and an improved wire quality with a lower gas and oxygen content can be obtained (Wincierz, 1975). The melt flows out of the casting furnace

6.5 Technology and Products of Copper-Based Alloys

into a cooled ingot mold (casting wheel), and is rolled afterwards in several rolling steps to an intermediate size of about 6-mm diameter. By this means a high-quality conducting wire of OF-Cu is produced in a controlled atmosphere with a high conductivity, low recrystallization temperature, and a smooth surface. Another wire fabrication process using OF-Cu is dip forming. A core wire is dipped into the oxgen-free melt; the wire diameter increases as a result of all-round crystallization, the emerging wire cools in an oxygen-free atmosphere or in a vacuum, and is then rolled and drawn continuously as usual (Diehl, 1991). By this process, wire products of pure copper, lowalloyed copper, and copper alloys of tin bronze, nickel-silver, and yellow and red brass are produced, which meet the highest application requirements. Standard and special wires are mainly used in all electrotechnological parts, for eye-glass frames, and in zip-fastener industries. Where good contacts are requested, copper wires are produced in the shape of wire-wraps and termipoint wires, e.g., tin bronze either bare or tin-plated whiskerproofed surfaces for connectors in computers and electronic circuits. Other kinds of connection are made using shaped-wire profiles of nickel-silver, e.g., for model railway tracks, zip-fastener wires, and optical wires with a special quality surface and protection (Diehl, 1991).

325

strength and a smooth surface locally, which is as good as that of soft-annealed copper. Die-forged pieces have a tight (pressed) microstructure in comparison to cast pieces and, accordingly, a high conductivity and good mechanical properties, especially in the highly deformed zones of the pieces. Highly conductive copper is deformed in neutral or reducing atmospheres at up to 790 °C, and oxygen-containing copper and brass at up to 900 °C. Die forging needs expensive dies, and is therefore suitable for mass production in the electrical, fitting fabrications, and vehicle industries, e.g., electrical contacts, terminals, and switches, thermostatic and water valves, and synchronizer rings and gearshift forks for mechanical transmissions (Diehl, 1991). Free-form forging is only economical when applied to small quantities. The starting material is in the form of extruded rods, although preforms are also needed for large units. The preform cast is heated homogeneously in furnaces at forging temperatures between 800 and 950 °C. Then it is deformed with simple tools in hammer mills or in hydraulic presses, e.g., stretching, spreading, punching, cutting, or bending. Free-form forging is not as precise as die forging (the admissible deviation of the dimensions is larger), but can produce bigger parts with high strength and good surface quality. 6.5.3 Copper Castings

6.5.2.5 Hot Forgings Forgings are deformed semiproducts of copper, which are produced in dies or free forms and then hot or cold deformed. Other die forgings are manufactured from copper by hot pressing followed by cold pressing. Sufficiently strained parts obtained by this cold pressing can have a higher

The two principal groups of unalloyed copper castings are distinguished according to their applications: cast copper mainly for nonelectrical applications (G-CuL35 with low conductivity, L indicates the electrical conductivity in m/Q mm 2 ), and conducting copper with high conductivity and purity (G-CuL45 and G-CuL50). For elec-

326

6 Copper-Based Alloys

Table 6-16. Casting (sand castings) of pure copper used as conducting materiala. Designation15 Alloying composition (wt.%)

Electrical conductivity

Heat Yield Tensile Fracture Brinell Applications conduc- strength strength elonhardness tivity gation at20°C (m/Qmm 2 ) (W/mK) (N/mm2) (N/mm2) (%)

G-CuL35

Cu98 (Zn,Sn)

>35

>178

45

170

25

42

cooling elements, sealing rings, fittings, etc.

G-CuL45

Cu99.8

>45

>314

40

150

25

40

G-CuL50

Cu99.9

>50

>348

40

150

25

40

parts for electrotechnology, switching elements, contact pieces, conductor parts, etc.

G-SCILL50

Cu99.9

>50

>348

40

150

25

40

a

Minimum values, DIN 17655; soldering.

b

G = sand-mold casting; L = electrical conductivity; S = welding and hard

trically conducting connections formed by welding or hard soldering, G-CuL50 is preferred (Table 6-16). Technical specifications for cast material with high temperature strength, fine grain size, and tightness generally require an additional 1-2% alloying element added to the pure copper. These normal copper castings (G-CuL35 or G-CuL50) still have a good conductivity of at least 35 or 50 m/ Q mm 2 (Table 6-16). The strength and hardness of copper castings are relatively low, nevertheless the castings have other important properties (Table 6-17) (i.e., high conductivity, good corrosion resistance, sufficient ductility, and easy alloying). Hence a number of alloying elements can be used, e.g., tin, zinc, aluminum, lead, nickel, or beryllium, which singly or in combination improve some of the properties of pure copper or satisfy the given requirements. These copper castings (Table 6-18) have, for example, a high static and dynamic strength, good high temperature strength, high hardness, improved wear resistance, excel-

lent corrosion resistance, good sliding behavior, and a partially improved castability. Copper castings have essentially three areas of application: as conducting materials, as bearing materials, and as construction materials for components subject to stress corrosion. 6.5.3.1 Conducting Materials Unalloyed and low-alloyed copper, which has a determined limiting electrical conductivity value, is used as conducting material (with a conductivity of 45 m/Q mm 2 (L45) and 50 m/Q mm 2 (L50); Table 6-11). The different forms are used chiefly for conducting components of all sorts, e.g., switches, conductors, contacts, and electrode supports. In addition, conducting Cu-Zn castings of higher strength are used, e.g., G-Cu-Znl5 with % = 15m/Q mm 2 or G-Cu-Zn40-Fe with x = 10 mm/Q mm 2 . These materials also have good thermal conductivity and are used in, e.g., armatures, seals, and cooling equipment.

6.5 Technology and Products of Copper-Based Alloys

327

Table 6-17. Physical properties of copper casting alloysa. Alloying group

Density

(g/cm3) Unalloyed copper castings Cu-Sn castings Cu-Sn-Zn castings Cu-Zn castings Cu-Zn special castings (brass) Cu-Al castings Cu-Pb-Sn castings Cu-Cr castings Cu-Ni castings Cu-Ni-Zn castings (nickel-silver) 1

8.9 8.6-8.7 8.7-8.8 8.5-8.6 8.2-8.6 7.5-7.6 8.7-9.5 8.9 8.9 8.9-9.0

Melting Electrical Thermal Linear temperature conductivity conductivity thermal range at 20 °C at 20 °C expansion 25-300 °C (°C) (m/Qmm 2 ) (W/m K) (10" 6 /K)

1083 820-1010 840-1045 880- 940 850- 900 1025-1050 845=1050 1075-1085 1105=1240 1010=1180

35-50 6.2-7.0 6.5-8.5 10-15 4-9.5

170-335 54- 59 56- 71 65- 90 34- 60

2-8 6-8.5 45 2.5 = 5.5 3-4

50- 60 54™ 71 295 29- 59 23- 33

17 18.5 18.0-18.8 19-20 18.5-20

Specific heat capacity

Elastic modulus

(J/g K)

(kN/mm2)

0.38 0.36 0.38 0.41 0.42

110 93.5-98 90- 93 93- 98 100-106

16-19 0.46 113-121 18.5-19.3 0.34-0.42 78.5- 85 18 0.38 110-130 15-16 0.38 123-145 17-19 0.42 128-137

According to ASTM and DIN standards.

6.5.3.2 Bearing Materials Bearing materials have to fulfill a lot of different requirements (Pfestorf et al., 1985), which will be referred to as 'sliding properties'. Many copper castings have good sliding properties, and therefore heterogeneous Cu-Sn, Cu-Sn-Zn, and Cu-PbSn castings are used to produce bearings, sliding strips, gears, sliding blocks, and sintered disks, running with oil or dry lubrication. The manufacture of these castings, which are often rotationally symmetrical, is achieved economically by centrifugal casting or extrusion. Important criteria for bearings are their wear and sliding properties; those of copper castings are excellent. Cu-Pb-Sn castings have good sliding and emergency run characteristics (G-Cu-PblO-Sn), as do CuSn castings (G-Cu-Snl2). The microstructure of Cu-Sn bearings consists of a soft oc-matrix with incorporated hard and fine dispersed eutectics (of 5-

and (3-phases). The high speed of solidification of the fine, dispersed eutectics during centrifugal casting or extrusion processes generates a higher volume fraction of eutectic in a finer distribution, so that the cast material is improved considerably compared to that made by sand casting. A lead addition is favorable for better emergency run behavior with mixed lubrication. The material then consists of a hard matrix with soft, dispersed particles, which, e.g., improve the embeddability. 6.5.3.3 Construction Materials In comparison to bearing materials, construction materials need to be hard and homogeneous, e.g., Cu-Zn or Cu-Al castings. For these applications the high strength through alloying and the good corrosion resistance of the homogeneous structure are important. Therefore, e.g., G-Cu-Zn37-Pb, G-Cu-Zn34-A12, GZ-, GK-Cu-Zn25-A15, and G-Cu-Alll-Ni are

328

6 Copper-Based Alloys

Table 6-18. Mechanical properties of copper casting alloysa. Alloying group b, material examples

Tensile strength,

Fracture elongation

(N/mm2)

(%)

260 280 260 280 280 270

12 14 10 5 7 18

80 90 80 90 85 70

Cu-Sn-Zn castings (red castings) G-Cu-SnlO-Zn 130 G-Cu-Sn7-Zn-Pb 120 90 G-Cu-Sn5-Zn-Pb 90 G-Cu-Sn2-Zn-Pb

260 240 220 210

15 15 16 18

75 65 60 60

8.7 8.8 8.7 8.7

bearings, gears, sliding strips and blocks, resistant against corrosion, erosion, and cavitation, armatures

Cu-Zn castings (brass) G-Cu-Znl5 70 70 G-Cu-Zn33-Pb 120 GD-Cu-Zn37-Pb GK-Cu-Zn37-Pb 90 GZ-Cu-Zn35-AH 200 GZ-Cu-Zn34-A12 260 480 GZ-Cu-Zn25-A15

170 180 280 280 500 620 750

25 12 4 20 18 14 5

45 45

8.6 8.5

70 120 150 190

construction material, plates, wires, lamps, gears, wheels, cranes, nuts

Cu-Al castings G-Cu-AUO-Fe G-Cu-A19-Ni G-Cu-Alll-Ni GK-Cu-A18-Mn

180 200 320 200

500 500 680 450

15 20 5 30

115 110 170 105

7.5 7.5 7.6 7.5

construction material, corrosion resistant, nuts, spindles, worms, gears, wheels

Cu-Pb castings G-Cu-Pb5-Sn G-Cu-PblO-Sn G-Cu-Pbl5-Sn

130 80 90

240 180 180

15 8 8

70 65 60

8.7 9.0 9.1

machining parts, bearings, pumps, bushes

Cu-Ni castings G-Cu-NilO G-Cu-Ni30

150 230

310 440

18 18

100 115

8.9 8.9

Unalloyed copper castings G-CuL35 45 G-CuL45 40 G-CuL50 40

ship construction, water installations. sludge fittings, pumps, stirring equipment

170 150 150

25 25 25

42 40 40

8.9

Yield strength, ^p0.2

(N/mm2) Cu-Sn castings (bronze) G-Cu-Snl2 G-Cu-Snl2-Ni G-Cu-Snl2-Pb GZ-Cu-Snl2-Pb GC-Cu-Snl2-Pb G-Cu-SnlO

a

140 160 140 150 140 130

Density

Brinell hardness, HB10/1000

Application fields

Q

(g/cm3)

75

8.6 8.6 ] ,

1

] i

8.7

bearings, gears, sliding strips and blocks, resistant against corrosion, erosion, and cavitation

8.7

o c O. J

8.6 8.6 8.2

electrical and thermal conductors, switches, contacts, short-circuit rings, electrodes supports, seals, cylinder heads

DIN 1705,1709,1714,1716,17655,and 17658, minimum values; b G = sand-mold casting, GK = chill casting, GZ = centrifugal casting, GC = continuous casting, GD = die casting, L = electrical conductivity value.

329

6.5 Technology and Products of Copper-Based Alloys

used (Table 6-18). Cu-Ni castings are used in ship construction, hydraulic engineering, the chemical industry, seawater desalting plants, fittings, pumps, and for applications requiring elevated impact behavior at low temperatures as in the refrigeration industry. Copper castings are produced by all kinds of die and pressure processes, similar to iron materials. Economical factors favor cast material, mainly produced in sand molds (abbreviation G) or chills (GK), or with centrifugal (GZ), extruded (GC), or pressure (GD) processes. The volumes, masses, inserts, and molds, however, are very different for products of 5 kg and 50 t. For molds with long lifetimes and precise measurements, Croning's process with mask, plaster-cast, or fine-casting processes into ceramic or metal molds (e.g., G-CuZnl5-Si4) are suitable. To obtain a homogeneous microstructure, uniform wall thickness, and optimal fiber sequence, the specific material and casting properties must be taken into account when the molds are constructed.

The roughness of the cast surface depends on the casting process. Typical values are: for sand castings, Ra = 18-80 jim, for castings in masks, Ra = 4-20 |im, and for precision castings, Ra = l-7 jim. The shrinkage of copper alloys is between 1.2 and 2.1%, depending on the process and the length of material (Weisner et al., 1982). Combined castings of liquid metal and a solid metal part are made by shrinking, rolling, casting, or clamping the liquid metal around the solid metal. In this way, a junction can be made by notches, grooves, or rough surfaces, but also by diffusion of the metals (e.g., copper and deposits of copper or iron) or by an intermediate layer (e.g., of silver, zinc, or tin), which increases the diffusion between both pieces. 6.5.4 Sintered Copper

Sintered pieces with good sliding behavior, but low wear, and in special forms can be produced from copper powder (Table 6-19).

Table 6-19. Properties of melted and sintered copper in various fabricated forms3 Fabricated forms

As cast Melted, forged, rolled, soft-annealed Melted, rolled, 80 % deformed Cold-pressed powder (600 MPa), sintered at 600 °C Cold-pressed powder (3000 MPa), sintered at 600 °C Powder hot-pressed (800 MPa) at 300 °C Powder hot-pressed (1500 MPa) at 400 °C Cold-pressed powder (300 MPa), sintered at 900 °C, 90% hot-deformed, annealed Cold-pressed powder (300 MPa), sintered at 900 °C, hot-forged, 50% cold-deformed 1

DKI (1982).

Tensile strength (N/mm2)

Elongation

Density

(%)

(g/cm3)

40-50 >50 100-110

150-195 195-235 430

15-25 >38 10-12

8.9 8.93 8.93

45 <40

140-155 -

3-4 2-3

7.6 6.9-7.0

100-110 120-150

245-295 295-345

10-20 -

8.9 8.9

40-50

195-245

<35

8.9-8.93

100-110

345-390

10-12

8.9-8.93

Hardness

330

6 Copper-Based Alloys

Copper powder, produced by electrolysis has dendritic particles, while copper powder produced by water spraying has spherical particles. These copper powders of different grain sizes can be pressed or rolled with other metal powders or metal oxides, for example, for sliding contacts. Another process is sintering, by which the pressed piece is strengthened during heattreatment below the melting temperature by diffusion, evaporation, and plastic deformation. Cu-Zn alloys with 10-30% Zn can also be hot pressed and sintered (ASM, 1979; DKI, 1982). In the form of sintered metal containing pores, the copper can be used for gas and liquid filters and for self-lubricating bearings, because the coherent pore network has a high storage volume for lubricants. Such filters are used to reduce the pressure, for homogeneous distribution of the gases, or for protection against back ignition of explosive gases. For self-lubricating bearings Cu-Sn powder alloys are mostly used, and the pores are filled with lubrications. These bearings are sintered or sprayed from 90 % copper powder and 10% tin powder or powder from other copper alloys. For coatings and materials with high friction coefficient, copper, tin or bronze powders are mixed and sintered with metal oxides or silicates. Sintered friction linings can be substituted for organic linings in brakes and clutches. Sintered linings do not pollute the atmosphere with carcinogenic dust, like that released by organic linings. The developed spray-sintering process can produce friction disks of sintered bronze (copper-tin) for the toughest application i.e. gear-boxes in ships (Diehl, 1991). Sintered disks are also used in automatic transmission clutches for buses and lorries, or for clutches in oil sumps of

heavy agricultural and construction site vehicles, or brakes in cranes and elevators. Their functions include high friction values, low wear, good thermal conductivity, and compact construction. Sintered bearings are generally produced from copper-tin (9-11% Sn) with or without graphite (0.2-2% C) (SINTB50, -B51, -C51). They are corrosion resistant and produce low noise. They are standard materials which are used in precision mechanics and combustion engines (Pfestorf et al., 1985). 6.5.5 Fabrication of Products by Non-Cutting Techniques

The good deformability of copper and its alloys is the basis of fabrication without cutting. Copper can be readily deformed with a low amount of lubrication. The low deformation forces, the simple formation of a lubricating film, and the good ductility of the oxide layers at the surface work together here. Hot deformation of unalloyed copper is usually done between 550 and 600 °C, during which the material is dynamically recystallized so that a lasting strengthening cannot be obtained. All hot deformation processes, e.g., rolling, extruding, and forging, can in fact be carried out over a wide temperature range. Copper is an ideal material for hot deformation. For copper alloys, the temperature of hot deformation is shifted mostly to higher temperatures, and higher forces, depending on the strength required. Cold deformation strengthens the material, which increases with progressive deformation, and progressively decreases any possible alteration (more dislocations in the microstructure). Cold deformation of copper is normally done at room temperature. When the total deformation en-

6.5 Technology and Products of Copper-Based Alloys

ergy is transformed into heat (fast deformation), the temperature of a heavily deformed workpiece can reach the recrystallization range. Consequently, the material is weakened. Rolling and drawing processes are used for the production of copper semiproducts (plates, strips, wires, profiles, rods, and pipes). For further fabrication of copper and copper alloys, different processes are used; e.g., cutting, deep drawing, stretching, pressing, rolling, and flow pressing into large or hollow molds. Products of copper and copper alloys can either be formed by solid deformation or by sheet deformation. Solid deformation (with reduction of the thickness of the slug) is carried out in the cold condition (cold rolling, drawing, upsetting, pressing, coining, flow pressing, die forming, and thread and profile rolling) or in the hot condition (hot rolling, forging, and extruding). For cold deformation, higher forces are normally required than for hot deformation, but a higher surface quality can be attained. Generally, those copper alloys with homogeneous a-structures are best cold deformed, while heterogeneous structures (e.g., a + P in Cu-Zn alloys) are more easily hot deformed if one phase (P) is softer at high temperatures. During sheet deformation, different processes (e.g., bending, deep drawing, stretch drawing, widening, and pressing) can often be used with only intermediate annealing, depending on the strength of the materials (Wieland, 1986). 6.5.6 Heat-Treatment and Annealing By soft annealing (softening), part or all of the work hardening can be removed, and so further deformation is feasible. Deformation grades with 25-90% (or more) area reduction can be obtained without intermediate annealing of f.c.c. oc-copper.

331

The minimum annealing temperature is about 200 °C for copper; soft annealing is usually done between 400 and 500 °C. If there are additions and alloying elements this temperature must be raised significantly, because of the higher recrystallization temperature of the alloys. With higher temperatures shorter annealing times are possible, but then more scale is formed and the grain size is larger. The recrystallization temperature of copper and copper alloys is normally between 500 and 600 °C depending on the previous cold deformation (Fig. 6-9). The stressed and deformed crystal structure changes into a new, unstressed structure. Here three steps are observed with increasing time and temperature: recovery, recrystallization, and grain growth (see Vol. 15 of this Series). The stress-free annealing temperature is low for copper and many copper alloys at -200-300 °C with 0.5 h annealing time. Herewith the structure is released and dislocations partly disappear (polygonization), but the grain size is not changed. Diffusion annealing is carried out at very high temperatures but below the melting temperature, e.g., about 700-900 °C for copper and brass (Fig. 6-19). By diffusion (atomic location change) it is possible for impurities to be completely dissolved in the solid solution and for the material to become softer with better conductivity. This condition is preferred for hot deformation (forging). With increasing temperature the strength is reduced and deformation is easier, but the formation of scale increases considerably, leading to material loss and a worse surface. For this reason, the usual hot deformation temperature is between 800 and 900 °C for copper and brass. A lower temperature results in a finer grain size. If a better surface quality is required,

332

6 Copper-Based Alloys

a lower deformation temperature is preferred (profiles, die forgings, and forgings). If a bright surface is desired, the copper material must be annealed in a furnace with a reducing (not oxidizing) atmosphere, e.g., soft annealing of copper wire in a continuous furnace. During annealing of oxygen-containing copper (above 560 °C), hydrogen must be removed from the reducing protective gases because of hydrogen embrittlement. The protective gas atmosphere must be neutral or slightly oxidizing (DKI, 1982). The cooling or quenching speed is not important for unalloyed copper, but must be determined for certain copper alloys. 6.5.7 Metal Cutting Tool wear is relatively high in the machining and cutting of copper. Copper and some soft copper alloys have a high tendency to smear and to paste over. The cutting tools need to have a high wear resistance and strength at their edges, therefore high-speed steels and hard metals are mostly used. Diamond tools do not have a tendency to paste over, but are only used for small parts and for fine machining. The cutting forces for copper are low in comparison to those for steel. The surface quality of soft copper depends greatly on the tool edge geometry and on chip formation. The chips are plastically deformed during cutting and become tough and hard, and they can cause problems while machining the soft copper. Short chips are formed with lead (Table 6-18), tellurium, sulfur, and selenium additions to the copper alloys, and hence such alloys can be machined with automatic machinery. In addition, heat-conducting lubricants are used, e.g., oil mixtures with and without additions of hard metal.

All metal-cutting processes are used for copper alloys, e.g., turning, drilling, planing, grinding, milling, sawing, and threading, but attention must be taken of the mode of chip formation of the soft copper material (DKI, 1982). 6.5.8 Recycling of Copper-Based Materials One aim of the development of material is to preserve the available primary raw material. Copper can be saved by better utilization of the raw materials, by energysaving production, and ultimately by environmentally conscious recycling in different ways. Scraps of one metal (or those that can be separated into various materials) can be put directly back into the material loop. High melting temperatures can burn out the organic components of the metal scrap, and as the slag swims on the metallic melt because of its low density it can easily be removed. The copper can be metallurgically reclaimed from the melt. Difficulties in recycling arise with aggregates and components which include a large number of different materials, e.g., vehicle bodies, household equipment, refrigeration equipment, industrial scrap, communication apparatus, electronic scrap, and polluted packing boxes. The amount of recycled copper materials (Fig. 6-6) is at present, in the technological countries, about 40 % of the total copper production (DKI, 1982; Weber, 1989; Metallgesellschaft, 1993). Recycling conserves the deposits of copper raw materials and reduces energy consumption. The production sequence from the raw material to the end product and copper production from ores consume about five times more energy than production starting from scrap. So new recycling techniques should be developed and used (Metallgesellschaft, 1993).

6.5 Technology and Products of Copper-Based Alloys

Copper is a material that can be used over and over again, as it is impossible to distinguish recycled copper from primary copper. This means that resources can be conserved and energy saved at the same time. The main process for copper recycling is electrorefining, which removes both base-metal and noble-metal impurities.

6.5.8.1 Classification The different kinds of scrap can be classified by their purity, their chemical composition, or by their origins: (1) Pure copper scrap can be remelted directly at production plants, or if it has higher levels of impurities it can be refined. (2) Most copper scrap, however, is alloy scrap such as brass, bronze, or nickelsilver. If it is very pure, this scrap can be directly melted down into alloy products. Much alloy scrap is also blown in a scrap converter to remove impurities such as zinc, tin, or lead, and then refined into high-purity copper by casting into anodes and electrorefming (Fig. 6-36). (3) Besides these types of scrap, which are clearly defined by their origin, there are many kinds of mixed scrap, such as copper scrap that is contaminated with plastics or organic materials, and is therefore processed in a converter. Scrap processing gives many intermediate products like flue dust, metal-rich slag, ash, or mattes if sulfide primary materials are charged together with the scrap. (4) In addition, more and more special materials and waste products are appearing, which require special processing routes, e.g., cable scrap, electronic scrap with its extremely complex composition, etching solutions from the circuit-board industry, electroplating sludges, or catalyst residues.

333

Many by-products of a smelter can be returned into an integrated process with copper scrap and primary copper to minimize residues and primary materials (Fig. 6-36). Among the types of copper scrap which are generated in large quantities are electronic scrap (Fig. 6-37), motor vehicle scrap (Fig. 6-38), and etching solutions from circuit-board production (Metallgesellschaft, 1993). A lot of problems exist in relation to recycling conditions and further development, e.g., component miniaturization, laminated compound materials, waste gas and waste solution treatment, energy and cost problems, cost and space for landfilling, low-emission pyrometallurgical processes (electric furnace), shredding and pretreatment techniques to improve the separation of plastics, copper, and other metals, and solvent extraction by ion exchange for hydrometallurgical processes. Also the development of processes for the utilization of the plastics content of the secondary materials is needed. Whereas copper is easily recycled without loss of value, there are problems with plastics and adherent organic materials (insulation). For many plastics/metal compounds, thermal treatment is the only method of separation. 6.5.8.2 Electrical Scrap Electrical and electronic scrap is widespread and consists mostly of circuit boards and recyclable copper. Electronic scrap can be reprocessed in different ways, e.g., selective leaching (noble metals), mechanical pretreatment (the main problem is the plastic content), incineration (problems: the removal of plastics, the formation of dioxin, and high off-gas rates), pyrolysis (plastics are decomposed at 400600 °C in a hermetically sealed furnace

334

6 Copper-Based Alloys Copper concentrate (30%Cu) Residues, sludge, slime

Silicon, oxygen, air Electrical scrap, slag, dust Copper cement (fine, Cu-rich)

i Blast furnace or flash smelting furnace

Sulfide sludge (from waste water treatment)

Copper matte (63 - 90 % Cu)

Silicon, oxygen, air

Bronze, brass, copper scrap (baled scrap, 40 - 90%Cu)

Matt scrap metal converter

Flue dusf Sulfuric acid Recycle slag

Blister copper (96-98 % Cu)

Natural gas

Copper granulates (from cable shredding > 90% Cu)

Flue dust Sulfuric acid Slag Zinc oxide, mixed tin

Anode furnace

Recycle copper slag

Anodes (98 -99.5%Cu)

Sulfuric acid

Nickel sulfates Copper electrolysis Anode slime (Se,Ag,Au,Pd,Pt)

T

Copper cathodes (99.99%Cu)

Fabrication Figure 6-36. Copper recycling. Integration of secondary copper materials into primary copper processing (Metallgesellschaft, 1993).

Glass, ceramics 2.7% Metals 60.5%

Remainder 7.8% Electronic scrap 11.5% Cables 4.4% Plastic materials 13.1%

Figure 6-37. Recycled materials. Distribution of material groups in a Siemens study of 2000001 of industrial electronic scrap (Kreft, 1993).

6.5 Technology and Products of Copper-Based Alloys

335

Automobiles

Dismounting of components

Components (radiator, alternator, battery)

Shredder

Shredder waste

Air classifier

Lightweight shredder waste

Iron removal magnet

Ferrous fraction

Nonferrous fraction

Cryogenic shredder

Iron-

Heavy media separation

— — Waste

M Nonferrous metals Materials Sorting """ (Al, Zn, Ni, etc.)

Iron removal magnet

Nonferrous metals

Copper alloys, copper

Copper smelter

with small amounts of flue gas), and processing in a converter (the plastics are destroyed, but the copper recycling is economical, limited capacity). More than 1.5 million tonnes of electrical scrap is produced per year in Germany (Kreft, 1993). Different types can be distinguished: about 400 000 t household appliances (~55%), 240 0001 consumer electronics. 260 000 t from the electrical engineering industry, 1800001 communications technology and the rest capital goods, e.g., laboratory, office, and medical equipment and plant (Handelsblatt, 1993). Very little electrical scrap consists of electronic circuits and components (~ 3 %). In total, about 800 000 t electronic scrap goes yearly in to landfill, without utilization.

Figure 6-38. Copper recycling. Recovery of copper materials from old automobiles (Metallgesellschaft, 1993).

Because of the order of magnitude of electronic scrap in Germany (1994), the producer and trader are obliged to take back and to utilize electrical and electronic equipment. With electrical scrap, the separation of material for reuse is difficult. In particular, scrap with high amounts (~30%) of plastics and nonmetals (light fluff) produces second class material for applications with lower requirements or frequently goes straight into landfill. A study by Siemens AG (Kreft, 1993) investigated the distribution of types of materials making up 200 000 t industrial electronic scrap that the company was forced to take back in 1993, and which could be reutilized. The major part of this consisted of metals (Fig. 6-37).

336

6 Copper-Based Alloys

6.5.8.3 Automotive Scrap

6.6.1 Corrosion-Resistant Materials

The nonferrous metal fraction of automotive scrap is often still separated by hand into the various alloys (Fig. 6-38). Copper from dismantled cars is recovered in various forms: as alloy scrap (such as brass or bronze in radiators), as copper granules from cable stripping, as copper/ iron scrap from the shredder, or as circuit boards from car electronics. These materials can be processed in blast or electric arc furnaces, scrap metal converters, or cast at anodes (Fig. 6-36).

Numerous applications of copper and copper alloys are based on its corrosion resistance. Copper is a nobler metal than others, and therefore does not go into solution so quickly as baser metals, like iron, zinc, and aluminum. For this reason, copper is very corrosion resistant in water (seawater, acids, and rain) and has a long lifetime. Due to its noble potential in liquids with hydrogen ions, copper is not corroded by water, aqueous solutions, or non-oxidizing acids. In neutral and alkaline aqueous liquids it is chemically resistant, and therefore it is used in water pipes, vessels in breweries, and for the storage of numerous liquids (Todt, 1961). Copper is corroded by oxidizing acids (e.g., HNO 3 , H 2 SO 3 ), which is why it cannot be used for the storage of foodstuffs, wines, or fruit juice if they contain such acids. As a result of this, poisonous copper acetate could be formed. An aqueous solution that contains bivalent copper ions (these will be transformed into monovalent ions) can oxidize metallic copper (DKI, 1982). In the open air, a protective patina is formed on the copper of basic sulfate (mostly in country air) and some copper carbonate (in city air) or sometimes copper chloride (in sea air), e.g., on roofs and gutters. The so-established oxidized layer of the patina (copper(i) oxide) changes from red-brown via dark brown to gray. The carbonic acid or sulfur dioxide necessary for its formation is taken out of the air. These kinds of protective layer are very compact, very corrosion resistant, and greatly delay any further corrosion. The formation of scale (corrosion layer of copper oxide) depends on the temperature. It starts at 250 °C and continues up to 1025 °C (red Cu 2 O layer); above 1025 °C

6.6 Copper Materials for Special Applications Copper and its alloys are multifariously used in modern technology quite apart from their applications as conductors, resistance, and contact materials, or even as construction materials with various other and very specialized properties, e.g., strength, sliding, cladding, and corrosion resistance (Covington etal., 1992). In the chemical industries, in pipelines, and in seawater aggregates unalloyed copper and some copper alloys are used because of their high corrosion resistance. As a compound material, copper powder is pressed and sintered together with other metal powders or metal oxides to give a heterogeneous material. Since ancient times copper has been used for jewellery and decoration, even for artistic objects or as coin material. A great number of future-orientated applications are now being investigated for copper and its alloys, as a result of which the development of the oldest metal is always continuing and can surely offer new marketing areas.

6.6 Copper Materials for Special Applications

only eopper(i) oxide (black CuO layer) is formed. At low temperatures, oxide formation is slow and can take years. The type and speed of copper oxide formation can be influenced considerably by alloying elements, e.g., aluminum and nickel reduce scale formation. Hydrogen or hydrogen-containing gases can seriously embrittle oxygen-containing copper at elevated temperatures (hydrogen embrittlement). SE-Cu is insensitive to hydrogen. Copper is stable, even in the ground, if it is not very acid or alkaline, or if potentials cannot be formed with other metals in aqueous solutions (DKI, 1982), The main materials for seawater applications are Cu-Ni alloys, in particular CuNilO-Fel-Mn. Also, recent developments (Diehl, 1991) with very small additions of arsenic to brass, while maintaining a single-phase a-structure, show that this brass then becomes resistant to corrosive tapwater. Dezincification, which usually retains the contour of brass parts (Cu-Zn39-Sn) predominantly in soft, chloride water, can be stopped by this means. Applications in seawater are mainly sea and coastal vessels, offshore oil and gas platforms, and their installations, desalination plants producing fresh water from seawater, coastal petroleum and petrochemical processing plants, and coastal electricity generating stations. The reason for its use is that Cu-NilO-Fel-Mn develops a protective film of three layers, resistant to both corrosion and marine biofouling (Pircher etal., 1985). Copper alloy claddings can be made by welding and are also used, like compound materials, for corrosion protection (Kocher, 1989). 6.6.2 Copper Coatings

Copper components can be covered with coatings or claddings of metal, organ-

337

ic plastic material, or enamel to improve the surface, e.g., to smooth, polish, insulate, or make it more resistant against corrosion. Metallic claddings are produced by rolling, forging, or welding onto the base (back) material, e.g., for bearings, for corrosion protection of chemical pipes, or for pressure vessels. Metallic coatings are applied on the copper by chemical processes, e.g., plunging, spraying, and sputtering, or by electrolytic separation or precipitation. Modern techniques utilize chemical vapor deposition (CVD) and physical vapor deposition (PVD) (Galvanotechnik, 1974). The coatings on copper or copper alloys are made of different materials, e.g., nickel or nickel-chromium (electrolytic deposition, for improved corrosion protection), chromium (molds for plastic materials), noble metals (for jewellery and copper decoration, electrical components, cutlery, arts, roofs), tin (electrolytic or galvanic deposits, layer thickness from 1 to 20 jam, good soldering and corrosive protection, wires, strips), lead (galvanic deposition, hot-water boilers), or tin-lead (for good soldering, integrated circuits). Organic coatings are layers made of lacquers and plastics, and used for insulation or coloring. They are deposited by plunging, sputtering, or spraying. Copper wires are laquered during drawing processes. Enamel coatings are produced on copper for decorative or artistic applications, sometimes also for architectural objects or components and wires for electrotechnology. The large difference in linear expansion between copper and enamel deposits does not cause any problems. By copper metallization, it is possible to vaporize very small films or interconnects (between 0.1 and 1 jam thickness), which are used as conductors in microelectronics (Aritaetal., 1994).

338

6 Copper-Based Alloys

6.6.3 Construction Materials Copper and copper alloys are not only used for electrically conducting material, produced from semiproducts and castings, but also as construction materials applied in the precision mechanics, machinery, and chemical industries. In these areas the efficient metal cutting (of the heterogeneous oc + J3 structure) of small construction pieces is important, just as is the good deformation (of the homogeneous oc-structure) of sheets, tubes, and profiles. Economical parts are also used here, which are sintered with copper and other materials. The conjunction of hot or cold forming with a uniform structure and good machinability gives a technical specification for high-tensile brasses (special brasses) that includes three preferences: medium to high tensile strength, improved corrosion resistance, and exceptional friction behavior. Frictional sliding of brass components is required in moving techniques and hence requires material of high strength. Various high-tensile brasses are used as bearing bushes, gear parts, and valve guides. Applications in engines and transmissions require copper alloys with special properties, e.g., wear resistance, thermal conductivity, ductility, and surface quality (Diehl, 1991). Cu-Zn alloys (brasses) are utilized in furniture, lamps, household articles, and jewellery, because of their gold-resembling color and appearance. Tubes, casting parts, plates, and wires are available in different colors and with different coatings. Brass tubes also bring comfort as heating tubes of different dimensions. Brass tubings transmit the heat in heat exchangers, which is transformed into warm air or liquid in, e.g., chemical plants. In mechanical engineering and chemical plants, copper and copper alloys are used

as heat exchangers, heating and cooling lines, copper pans, pipelines, soldering irons, solders, for metal sealings of pressure vessels, sealings for combustion engines, and also as cladding material (rolling, melting, explosion cladding) and protection against corrosion. In civil and construction engineering, many established applications are in roofs of buildings, gutters, eaves, downpipes, and metal fittings. In electro technology nearly 50 % of the total copper consumption of the world (Fig. 6-5) is used for conducting applications, e.g., wires, cables, strips, profiles, casting pieces, electrodes, and contact parts. 6.6.4 Mechanical Composite Materials with Copper In cases where high and contrary requirements are necessary for a material, which cannot be fulfilled by only one material, the diversity of composite materials gives different solutions and sometimes brings with it astonishing possibilities (Table 6-20). There are numerous examTable 6-20. Composite materials with copper for electrical applications3. Material

Applications

Cu-W, Cu/WC

contacts, sintered disks carbon brushes, sliding contacts system supports, semiconductor carriers compound wire, cover wire thermobimetal superconductor metal core substrate, circuit boards hybrid electronics, energy electronics

Cu-C Cu-Co Cu-Fe, Cu/Pb Cu/Cu-Ni Cu/Nb3SN, Cu/NbTi Cu/Invar/Cu Direct copper bonding (DCB) substrate Arpaci and Bode (1992).

6.6 Copper Materials for Special Applications

339

Brass . Cu-Zn 31-Si a-phase

Copper "a-phase

-Steel back

pies, for instance, if under thermal strain some different lengths are required (bimetals, claddings), or good electrical and thermal conductivity as well as high wear resistance are required (sliding contacts, brake disks). An example of a mechanical composite material of brass CuZn31-Si (oc-phase) with copper (a-phase), and steel is shown in Figure 6-39, used as bearing material. Forging is also selected for manufacturing bonded components of different conductive materials. Contact under high voltages is, for example, possible with forged composites of copper alloys (CuCd, Cu-Cr, and Cu-Ag) and aluminum. They have high conductivity, mechanical properties, surface quality, and resistance against corrosion in a desired combination and are used as electrical contacts, terminals, clamps, and switches. Another practice using combinations of various material properties is the incorporation of fine particles into the copper matrix (e.g., aluminum oxide particles in copper), by which the strength of copper alloys can be enhanced considerably (dispersion hardening). Dispersion-hardened copper alloys are produced for highly conductive contacts or for electrodes of resistance press welding, in which very small particles of metal oxides (e.g., A12O3,

Figure 6-39. Composite material of a-brass Cu-Zn31-Si with an intermediate layer of copper deposited on a steel back.

TiO2) are incorporated, dispersed in a mass of about 0.1%. These even remain dispersed inside the copper on heating above 400 °C, and their improved hardness is retained until close to the melting temperature. These dispersion-hardened alloys combine a low reduction of conductivity with good mechanical properties. On the other hand, the strength of precipitation-hardened Cu-Cr and Cu-Zr alloys is lost above 400 °C, because the particles are dissolved. On pressing and sintering, the metal powder and the copper powder, which is produced by electrolysis, are combined between 600 and 900 °C and are condensed with a reduced volume. Fine-grained SECu powder can be sintered with graphite powder (60-75% Cu) to form contact brushed or sliding contacts of carbon. Sintered W-Cu contacts have a high strength against burn out (Schreiner, 1964). Cu-Pb contacts (90% Cu) can be contacted by welding and are used for small switches. Sintered metal bearings are produced from copper or copper alloy powders, which are made by spraying or by electrolytic precipitation, e.g., tin bronze, brass, or iron powder. After sintering and calibration, the pores of these bearing brushes are filled with lubricating oil. Likewise, large amounts of copper powder

340

6 Copper-Based Alloys

are used as friction materials (brake linings). Inside the bronze (Cu-Sn-Pb) matrix of copper and tin powders hard metal oxides and silicates are incorporated and sintered, the bronze matrix providing the strength and conducting away the frictional heat. Electrolytic copper layers on steel, zinc, or aluminum confer good protection against corrosion and have many decorative and functional applications. Cu-Ni-Cr layers are frequently electrolytically deposited on the base material and polished to brightness afterwards. The copper layers or coatings have different combined properties, in particular electrical conductivity, corrosion resistance, the repair of worn metals, lubrication, and surface quality (Dettner and Elze, 1974). 6.6.5 Copper Alloy Bearing Materials A distinction is made between rollingelement bearings and sliding or plain bearings, depending on the load and the kind of friction. The most common material for ball and roller bearings is steel, while nonferrous metals are used for sliding bearings. These alloys have a heterogenous microstructure. These heterogeneous bearing materials can be classified into three groups: (1) materials where the microscopically visible particles are softer than the matrix (e.g., lead-containing tin bronzes and other metal alloys), (2) materials where the particles are harder than the matrix (e.g., aluminum alloys and the special brass Cu-Zn40-A12), and (3) materials with both hard and soft particles in the matrix (e.g., aluminum alloys) (Pfestorf etal. 1985). Another distinction is made between the metallic materials for sliding bearings, as outlined in the following sections.

6.6.5.1 Solid Sliding Bearings Solid sliding bearings consist of only one material. They are machined from castings, cold-drawn tubes, or roll-formed strips, using Cu-Pb, Cu-Sn, Cu-Zn, or CuAl alloys. • G-Cu-Pb20-Sn5, G-Cu-Pbl-Sn8, G-CuPblO-SnlO, and G-Cu-Pb9-Sn5 (G indicates casting) are castings for solid and multilayer bearings, for heavy loads and high speeds. • G-Cu-SnlO-P, G-Cu-Snl2-Pb2, G-CuSnlO-Zn2, and G-Cu-Sn5-Zn5-Pb5 are casting alloys for high loads and are seawater corrosion resistant. • G-Cu-Zn35-Al and G-Cu-Zn25-A15 are casting alloys for solid bearings with moderate sliding. • G-Cu-AUO-Fe, G-Cu-Alll-Ni, and GCu-A110-Fe5-Ni5 are castings for high stress and wear, and corrosion resistance. • Cu-Sn8-P and Cu-Sn4-Pb4-Zn4 are wrought alloys for solid bearings for high loads, moderate speed, and wear and corrosion resistance. • Cu-Zn31-Sil, Cu-Zn37-Mn2-A12-Si and Cu-Zn40-A12 are also wrought alloys for solid bearings for high loads and wear resistance. • Cu-A19-Fe4-Ni4, Ci-A110-Fe3-Mn2, and Cu-Alll-Ni6-Fe5 are wrought materials for solid bearings for heavy loads and excellent resistance against corrosion and wear. 6.6.5.2 Multilayer Metallic Materials Multilayer metallic materials for thinwalled sliding bearings are made from two or more layers of different compositions. The preferred manufacturing process includes casting, sintering, cladding by roll bonding, and galvanic deposition of the composite.

6.6 Copper Materials for Special Applications

341

Overlay (Pb-Sn10-Cu2) Ni-barrier

!u Pb22 Sn

Steel

Figure 6-40. Micro structure of a multilayer bearing (Cu-Pb22-Sn) with overlay. The overlay of PbSnlO-Cu2 is deposited on the steel-backed lead bronze (Pfestorf et al, 1985); inset enlarged by a factor of 5 compared to the main micrograph.

200pm

Steel-backed Cu-Pb alloys, which often have an additional overlay, are the most common. The very soft, very thin overlay, e.g., Pb-SnlO-Cu2 (Fig. 6-40), withstands the mechanical stress and serves as the sliding surface. The Cu-Pb22-Sn ensures that the bearing works, even if the overlay is subject to intensive wear. A very thin layer of nickel serves as a diffusion barrier between the overlay and the Cu-Pb22-Sn. Cu-Pb alloys containing tin have good emergency-run characteristics and embeddability, and allow a high sliding velocity. The higher the tin content, the higher the hardness, the load carrying capacity, and the corrosion resistance (Pfestorf et al. 1985). • Cu-PblO-SnlO and Cu-Pbl7-Sn5 are multilayer alloys with high fatigue strength and limited emergency-run behavior. • Cu-Pb24-Sn4, Cu-Pb24-Snl, and CuPb30 are multilayer materials with moderate fatigue strength and good emergency-run characteristics. In combustion engines with overlaid multilayer bearings of Cu-Pb24-Sn, specific loads of up to 70 N/mm 2 and slid-

ing velocities of up to 20 m/s are encountered. 6.6.5.3 Sintered Bearings Sintered bearings generally consist of only one alloy, which is often based on iron or copper. They are manufactured to achieve a defined level of porosity (1525%). • SINT-B50 (P4013 in ISO 5755) consists o f 9—11 % Sn, remainder copper, has a density £ = 6.8-7 g/cm3, a hardness, HB = 30, a good dynamic load capacity of 5 N/mm 2 at 1 m/s, and is the standard material. • SINT-B51 (P4022), and SINT-C51 (P4023) are tin bronzes with 9-11 % Sn, 0.2-2% C, (graphite), £ = 6.6-7.5 g/ cm3, and HB = 25; they have a low noise level and are used for precision mechanics and automobiles (Fedorchenko, 1983). 6.6.6 Decorative Materials Copper, tin bronze, and other copper alloys have a long history. In several cultural centers (Near East, Egypt, Aegean

342

6 Copper-Based Alloys

Sea, Middle East, and China), unalloyed copper objects have been found dating from more than 3000 years B.C., mostly household articles, jewellery, and weapons. Since the Bronze Age copper has been produced and made into jewellery, weapons, coins, tools, and die castings. Also the first higher strength copper alloys in bronze (Cu-Sn alloys) were used and copper winning in hearth furnaces was applied. Cu-Sn alloys (tin bronzes) gave their name to the Bronze Age and have been used since then up to the present day for jewellery and works of art. Cu-Sn-Zn alloys (red cast) are used, because of their good deformability and excellent corrosion resistance, for sculptures (e.g., Marcus Aurelius in Rome) or for sliding bearings and pump casings. Cu-Ni-Zn alloys (nickel-silvers) are used in manifold applications because of their silver-white color, high strength, and good corrosion resistance; applications include jewellery, arts and crafts, and springs for the electrical industry. Cu-Ni alloys are highly corrosion resistant and are applied therefore and, by reasons of their color, for coins, e.g., Cu-Ni25. The renaissance of Cu-Zn alloys (brass look) has been reflected in their utilization in furniture and sanitary equipment, lamps and other small commodities, as well as jewellery. Durable and transparent coatings preserve the beauty of the gold-resembling brass of various semiproducts and products for long periods.

6.6.7 Copper Metallization The use and manufacture of integrated circuits (ICs), e.g., DRAM (dynamic random access memory), SRAM (static random access memory), and EEPROM (electrically erasable and programmable read-

only memory) devices, and logic devices with high circuit speed, high packing density, and low-power dissipation, requires a reduction of the feature sizes in ultralargescale integration (ULSI) structures (Li et al. 1994). Here copper layers and deposits are applied in miniaturized interconnects of semiconducting chips. Interconnect feature sizes can be miniaturized by reducing the wiring pitch and adding more metallization levels. This increases the wiring resistance and capacitance, which increases the delay time and generates more cross-talk noise. Chemical vapor deposition (CVD) of copper films is now carried out by pyrolytic decomposition of copper(n) and copper©; the films produced have a special crystal structure. A wide variety of copper complexes exist for the deposition of pure copper films (0.5 jim thickness) with low electrical resistivities (ge = 2.0 JUQ cm). The choice of copper instead of aluminum is based on the new requirements of the manufacturing process. Complete, defect-free filling of vertical interconnects and the deposition of electrically uniform copper films (with roughly 3:1 aspect ratio) are problems, but can now be achieved (Arita etal., 1994). The resistivity of the metal interconnects may limit device performance in multilevel thin-film structures. For IC metallization, feature sizes of 0.25 (urn or less are required for the design of low-resistance metals, such as gold, copper, or silver. The future of new copper technology for ULSI interconnects with high quality copper, defectfree filling, and low cost has yet to be determined. Electroless copper metallization has already been utilized for semiconductors, based on the grain growth of deepsubmicron circuits. Hence copper will play a more important role in the semiconductor industry in the near future.

6.6 Copper Materials for Special Applications

6.6.8 Shape Memory Alloys Shape memory alloys vary their shape (volume) according to the temperature by undergoing a martensitic transformation with a determined sequence of stacking faults, as a result of stress and deformation variations. This phenomenon can be observed in a (3-Cu mixed crystal of brass e.g., Cu-Zn-Al, after an appropriate treatment. This results in pseudo-elasticity with a rubber-like behavior (a small change in stress results in high extension). The progress of stress and extension is reversible, but uses different routes, which can be remembered by the material. This behavior can be obtained after suitable treatment provided that the reversible transition of the body-centered cubic structure converts into another regular structure (e.g., martensite, P'-phase). Shape-memory alloys are used with copper in medical equipment (endoscopes) and the electrical industry (switching elements, storage, bending strips, thermobimetals) (Tautzenberger and Stockel, 1986; Grafen, 1991). 6.6.9 Other Applications of Copper 6.6.9.1 Copper as an Alloying Element Copper is an important alloying element for different metals. Copper is alloyed with cast iron, steel, aluminum, nickel, zinc, gold, and some other elements, to improve certain characteristic properties (e.g., strength, color, resistance, and conductivity), and to produce special semiproducts. 6.6.9.2 Chemical Compounds with Copper Copper has manifold applications in chemical compounds and copper salts, which have technical importance. Copper© oxide (Cu2O) is produced by oxidation of fine, dispersed copper, and has a

343

density Q of 5.8-6.2 g/cm3. It is used as paint for ships, as a pesticide, for the production of semiconductors, and for coloring glasses and enamel (yellow to redbrown). Copper(n) oxide (CuO) is formed as a black powder by oxidation through heating metallic copper in air and has a density Q of 6.3-6.4 g/cm3. It is used as a catalyst in various chemical processes, for coloring of glasses, and in brake linings. Above 100 °C, it changes to copper(i) oxide 4 CuO

(6-2)

Copper(n) sulfate (CuSO4) is an important copper salt (copper vitriol), has g = 2.3 g/cm3, and is used for the production of mineral colors, for plant protection, for the conservation of wood, and in galvanic baths. Technologically it is produced by dissolving copper chips in sulfuric acid in air. Copper(i) chloride (CuCl) is produced by dissolving copper chips in warm hydrochloric acid with an addition of potassium chlorate. It is a white, crystalline powder with a density of 3.5 g/cm3. It is used for the analysis of gases, and is oxidized in humid air. Copper(n) chloride (CuCl2) is formed together with hydrochloric acid from the original CuCl, and has a density of 2.5 g/cm3. It is used in photography and in petroleum refining, and in the oxidized condition to combat fungus in vineyards. Copper nitrate [Cu(NO3)2] is used for coloring copper, brass, and steel, and is produced by the dissolution of CuO in nitric acid (HN0 3 ). Copper carbonate [CuCO 3 -Cu(OH) 2 ] is used as a color in painting (e.g., azure blue), on china glazes, and in the paper industry.

344

6 Copper-Based Alloys

6.6.9.3 Copper in Medicine Copper additives are used in food (Grunau, 1970) and medicine. As a trace element in various plants and herbs, it is healthy and essential for life, e.g., in the formation of red-blood corpuscles, improved cellular respiration, and defence against infection (Leibold, 1994).

6.7 Directions for Future Applications of Copper An additional increase in copper production in the next two decades will be focused on new products, which are exclusively or partially made of copper, and at present already have industrial applications or signify new directions for future applications (Incra, 1989). Some of these new inventions and necessary investigations of copper alloys for the electrical industry are summarized as follows: • Modules for generating energy in industries with high thermal efficiency and low emissions of waste gases, which will be mounted completely during fabrication, and which are small units for gas firing plants (Arpaci and Bode, 1992). • New materials for circuits, coppermolybdenum foils with good electrical and thermal properties, high strength, and nonmagnetic behavior, as conductors. • Hybrid copper cables, electronic strips of Cu-Ni-Si alloys (Ruchel, 1989). • Copper structures inside stationary, large batteries, energy and supply to networks, or emergency current. • Protection of space against electromagnetic fields. • Electrical, diesel-electrical, or hybrid processes for powering vehicles and boats with novel storage batteries.

• Microwave equipment for generating heat (Arpacci and Bode, 1992). • Combinations with superconducting material for low-current circuits, conductors with a sputtered, cladd or drawn copper matrix. • Shape memory alloys based on copper. • Applications involving the superplasticity of copper materials, which have not yet been fully investigated (Zeiger, 1986). • The development of new materials for casting molds, casting wires, and casting cylinders (Gehrhardt, 1990). • Optimized fabrication processes for extruding copper hollow profiles (Gehrhardt, 1990). • Dispersion-hardened copper with optimized price-product relation and opening of new market, based on copper and A12O3. These examples show directly that copper can continue to be fully used in the technological future. Copper and its alloys have been and will always be modern materials. The old metallic material copper has accompanied the progress of techniques for the last 6000 years and will continue to do so.

6.8 References Arita, Y, Awaya, N., Ohno, K., Sato, M. (1994),

MRS Bull 19(8), 68. Armstrong Smith, G. (1970), in: Int. Conf. on Copper, Amsterdam (September 1970): London: Institute of Metals. Arpaci, E., Bode, A. (1992) Metall 46, 22. ASM (1979), Metals Handbook, 9th ed. Metals Park, OH: American Society for Metals (ASM), Vol. 2, pp. 237-490. ASTM (1991), Annual Book of ASTM Standards, Copper and Copper Alloys. Philadelphia, PA: American Society for Testing and Materials, Vol. 02.01. Bargel, H.J., Schulze, G. (1994), Werkstoffkunde, 6th ed. Diisseldorf: VDI-Verlag, pp. 256-266.

6.8 References

Beitz, W., Kiittner, K.-H. (Eds.) (1987), Dubbel Taschenbuchfur den Maschinenbau, 16th ed. Berlin: Springer, pp. E45-E51, E78-E85. CDA (1982), Aluminium Bronze Alloys. London: Copper Development Association, Technical Data No. 82. CDA (1990), Standard Designations for Copper and Copper Alloys. New York: Copper Development Association. CIDEC (1972), Copper-Nickel Alloys. Geneva: Conseil International pour le Developpement du Cuivre, Data Sheets. Covington, M. W, Church, N. L., Dresher, W. H., Cooper, B. R. (1992), in: Encyclopedia of Applied Physics, Vol. 4: Trigg, G. L., Immergut, E. M. (Eds.). New York: VCH, pp. 267-276. Dettner, W., Elze, J. (1974), Handbuch der Galvanotechnik, Munich: Hanser Verlag. Diehl (Ed.) (1991), Non-Ferrous Metal Products. Rothenbach: Diehl. Dies, K. (1967), Kupfer und Kupferlegierungen in der Technik. Berlin: Springer. DIN (1991), DIN-Taschenbuch, Vol. 26, Kupfer und Kupferknetlegierungen, Normen, 4th ed. Berlin: Beuth. DKI (1966), Kupfer-Zink-Legierungen. Berlin: Deutsches Kupfer-Institut. DKI (1976), Niedriglegierte Kupferlegierungen. Berlin: Deutsches Kupfer-Institut. DKI (1982), Kupfer, 2nd ed. Berlin: Deutsches Kupfer-Institut. DKI (1991a), Kupfer-Zink-Legierungen. Berlin: Deutsches Kupfer-Institut, Informationsdruck 1.5. DKI (1991b), Kupfer-Aluminium-Legierungen. Berlin: Deutsches Kupfer-Institut, Informationsdruck 1.6. DKI (1991 c), Kupfer-Nickel-Zink-Legierungen. Berlin: Deutsches Kupfer-Institut, Informationsdruck 1.13. DKI (1992 a), Kupfer-Nickel-Legierungen. Berlin: Deutsches Kupfer-Institut, Informationsdruck 1.14. DKI (1992 b), Kupfer-Zinn-Knetlegierungen. Berlin: Deutsches Kupfer-Institut, Informationsdruck 1.15. Fedorchenko, I. M. (1983), Powder Metall. Int. 15, 126. Fischer, H. (1978), Werkstoffe in der Elektrotechnik, Munich: Hanser Verlag. Galvanotechnik-Handbuch (191 A), 3rd ed. Berlin: Verlag Technik. Gehrhardt, J. W. (1990), Metall 44, 1076. Grafen, H. (Ed.) (1991), Lexikon Werks toff technik. Diisseldorf: VDI-Verlag, pp. 590-593. Grunau, E. B. (1970), Metall 2, 191. Guillery, P., Hezel, R., Reppich, B. (1991), Werkstoffkunde fur die Elektrotechnik, 6th ed. Braunschweig, Germany: Vieweg, pp. 70-86. Handelsblatt (1993), Das Elektronikschrott-Gebirge wa'chst. Diisseldorf: Verband Deutscher Maschi-

345

nen- und Anlagenbau (VDMA), Zentralverband Elektrotechnik- und Elektronikindustrie (ZVEI), Handelsblatt, 14.4.1993. Hansen, M., Anderko, K. (1958), Constitution of Binary Alloys, 2nd ed. New York: McGraw-Hill. Heller, W. (1976), Copper-Beryllium Alloys for Technical Applications. Geneva: European Organization for Nuclear Research, CERN 76-01. Incra (1989), Technology for the Copper Industry. Incra-Report No. 424. Jung, H., Matucha, K. H. (1993), in: Umformtechnik, Plastomechanik und Werkstoffkunde: Dahl, W, Kopp, R., Pawelski, O. (Eds.). Berlin: Springer, pp. 785-791. Kocher, R. (1989), "SchweiBen von Kupferwerkstoffen und Verbundwerkstoffen mit Plattierungsauflagen aus Kupferwerkstoffen", in: Jahrbuch Schweifitechnik '90. Diisseldorf: Deutscher Verlag fur SchweiBtechnik (DVS), pp. 35-48. Kreft, H. (1993), Siemens-Z. (5), 5. Leibold, G. (1994), Naturheilkunde. Niedernhausen/ Taunus, Germany: Bassermann Verlagsbuchhandlung, p. 78. Li, I, Seidel, T. E., Mayer, J. W. (1994), MRS Bull.

19(8), 15. Metallgesellschaft (1993), Die Welt der Metalle, Kupfer (The World of Metals, Copper). Frankfurt: Metallgesellschaft. Pawlek, E, Reichel, K. (1956), Z. Metallkd. 47, 347. Pfestorf, H., Weiss, F., Matucha, K. H., Wincierz, P. (1985), in: Ullmann's Encyclopedia of Industrial Chemistry. Weinheim: VCH, Vol. A3, pp. 399-420. Pircher, H., Ruhland, B., Sussek, G. (1985), Use of Copper-Nickel Cladding on Ship and Boat Hulls. New York: Copper Development Association (CDA). Richter, H. (1993), in: Umformtechnik, Plastomechanik und Werkstoffkunde: Dahl, W., Kopp, R., Pawelski, O. (Eds.). Berlin: Springer, pp. 717-721. Ruchel, P. (1989), Metall 43, 11. Ruge, J. (Ed.) (1991), Handbuch der Schweifitechnik, Band 1 Werkstoffe; Band 2 Verfahren und Fertigung, 3rd ed. Berlin: Springer. Schimpke, P., Schropp, H., Konig, R. (Eds.) (1977), Technologic der Maschinenbaustoffe, 18th ed. Stuttgart: Hirzel, pp, 120-144. Schreiner, H. (1964), Pulvermetallurgie elektrischer Kontakte. Berlin: Springer. Steffens, H. D., Sievers, E. R. (1990), Schweifien Schneiden 42, 501. Tautzenberger, P., Stockel, D. (1986), ZWF, Z. Wirtsch. Fertigung 81, 703. Todt, F. (1961), Korrosion und Korrosionsschutz, 2nd ed. Berlin: Walter de Gruyter. VoBkuhler, H. (1955), Z. Metallkd. 46, 525. Wallbaum, H. J. (1964), in: Landolt-Bornstein Zahlenwerte und Funktionen, IV. Band, Bandteil 2 b. Berlin: Springer, pp. 639-890. Weber, A. (Ed.) (1989), Neue Werkstoffe. Diisseldorf: VDI-Verlag, pp. 162-165.

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Weisner, E., Biichen, W., Kleinau, M., Bach, H. G. (1982), Gufi aus Kupfer und Kupferlegierungen. Diisseldorf: Gesamtverband Deutscher MetallgieBereien (DGM). (VDG). Wieland-Buch (1986), Kupferwerkstoffe, 5th ed. Ulm: Wieland-Werke. Wincierz, P. (1975), Z. Metallkd. 66, 235. Zeiger, H. (1986), Z. Werkstofftech. 17, 75.

General Reading ASM (1979), Metals Handbook, 9th ed. Metals Park, OH: American Society for Metals (ASM), Vol. 2, pp. 237-490. Bargel, H. I , Schulze, G. (1994), Werkstoffkunde, 6th ed. Diisseldorf: VDI-Verlag, pp. 256-266.

Beitz, W., Kiittner, K.-H. (Eds.) (1987), Dubbel Taschenbuch fu'r den Maschinenbau, 16th ed. Berlin: Springer, pp. E45-E51, E78-E85. Dies, K. (1967), Kuper und Kupferlegierungen in der Technik, Berlin: Springer. DIN (1991), DIN-Taschenbuch, Vol. 26, Kupfer u. Kupferknetlegierungen, Normen, 4th ed., Berlin: Beuth. DKI (1982), Kupfer, 2nd ed. Berlin: Deutsches Kupfer-Institut. Grafen, H. (Ed.) (1991), Lexikon Werks toff technik. Diisseldorf: VDI-Verlag, pp. 590-593. Guillery, P., Hezel, R., Reppich, B. (1991), Werkstoffkunde fur die Elektrotechnik, 6th ed. Braunschweig, Germany: Vieweg, pp. 70-86. Metallgesellschaft (Ed.) (1993), Die Welt der Metalle, Kupfer (The World of Metals, Copper). Frankfurt: Metallgesellschaft.

7 Nickel-Based Alloys Ulrich Heubner Krupp VDM GmbH, Werdohl, Germany

List of 7.1 7.1.1 7.1.2 7.1.3 7.1.4 7.2 7.2.1 7.2.1.1 7.2.1.2 7.2.2 7.2.2.1 7.2.2.2 7.2.3 7.2.3.1 7.2.3.2 7.2.3.3 7.2.3.4 7.3 7.4

Symbols and Abbreviations Main Fields and Requirements of Applications General Survey Wet Corrosive Applications High-Temperature Applications Electrotechnical Applications Composition, Structure, Properties, and Behavior Nickel-Based Alloys for Wet Corrosive Applications Composition, Structure, and Resulting Mechanical Properties Composition, Corrosion Behavior, and Resulting Special Uses Nickel-Based Alloys for High-Temperature Applications Composition, Structure, Mechanical Properties, and Resulting Applications Corrosion Behavior, Composition, and Resulting Special Uses Nickel-Based Alloys for Electrotechnical Applications Heating Resistance Alloys Spark Plug Electrode Alloys Soft Magnetic Alloys Controlled Expansion Alloys Concluding Remarks References

Materials Science and Technology Copyright © WILEY-VCH Verlag GmbH & Co KGaA. All rights reserved.

348 349 349 349 350 350 351 351 351 361 368 368 378 382 382 385 385 392 394 395

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7 Nickel-Based Alloys

List of Symbols and Abbreviations

Xt

activity of carbon tensile elongation induction remanence saturation induction magnetic field strength component saturation polarization magnetocrystalline anisotropy constant uniaxial anisotropy energy level of d-orbital average energy level of d-orbital rupture strength yield strength in terms of stress for 0.2% strain temperature Curie temperature take-out temperature molar fraction of component i

a X Xs fi4

coefficient of thermal expansion magnetostrictive anisotropy constant average magnetostrictive anisotropy constant initial permeability

jnmax

maximum permeability

ac AOD cct cpt dc HV IP LC mpy PWR RE

alternating current argon oxygen decarburization critical crevice corrosion temperature critical pitting temperature direct current Vickers hardness TEM intergranular penetration low carbon TTP TTS mils per year pressurized water reactor VCM VOD rare earth element

A5 B Br Bs H i Js Ki

Ku Md Md Rp02 T Tc TE

transmission electron microscopy time-temperature-precipitation time-temperature-sensitization vinyl chloride monomer vacuum oxygen decarburization

Trademarks Konstantan, Nicorros, Nicrofer, Nimofer, Cronifer, Cronix, Magnifer, and Pernifer are registered trademarks of Krupp VDM GmbH. Other trademarks referring to nickel-based alloys are addressed in Vol. 7, Chap. 14 of this Series.

7.1 Main Fields and Requirements of Applications

1.1 Main Fields and Requirements of Applications 7.1.1 General Survey

Most of the primary nickel produced in the western world is consumed in the manufacture of steels, predominantly stainless steels. During the three-year period 19901992 about 72% of the primary nickel was used for steel making. Only about 12% was consumed for making nickel-based alloys. About 9% served for electroplating and miscellaneous purposes. Approximately 2 % was needed for making nonferrous alloys other than those based on nickel, especially copper alloys. The total production of nickel-based alloys in the western world during the period 1990-1992 can be estimated as about 145000 tonnes/ year. This figure includes alloys with at least 30 mass% nickel, this is different from the German standard (DIN 17 600, 1969) definition according to which an alloy is defined to be nickel-based if nickel constitutes the biggest single mass percentage fraction. The most prominent features for the application of nickel-based alloys in today's industrial world are their high resistance to attack by aggressive corrosive media, both wet corrosive and high-temperature gaseous, their high creep strength at elevated temperatures, and certain physical properties. Each of these features are not necessarily universal characteristics of nickelbased alloys, but are related to special alloy systems. Therefore the main fields of application of nickel-based alloys, roughly estimated as about 30 % of their total consumption, are in the chemical process industry and in wet corrosive applications for electric energy generation, including feed water heaters and steam generation tubing, as well as pollution control equip-

349

ment, especially for flue gas desulfurization. Another percentage, which is roughly estimated as about 10%, is used in marine and offshore engineering. Important hightemperature applications, estimated as about 20 % of the total alloy consumption, are in the hot sections of airborne and stationary gas turbines and in industrial furnaces. Soft magnetic and expansion controlled alloys for electrotechnical components constitute another estimated share of about 20 % of the total nickel-based alloy consumption. The remainder is essentially heatiAg resistance elements, automotive applications (e.g., spark plug electrodes), and coinage. Most nickel-based alloys are used as wrought materials. Foundry alloys constitute only a minor fraction of the nickelbased alloys used in the chemical process industry and in electric energy generation. They are therefore not within the scope of this text. In contrast, foundry alloys, mostly investment cast, are of great importance for manufacturing blades and vanes for airborne gas turbines. These are the subject of Vol. 7, Chap. 14 of this Series. 7.1.2 Wet Corrosive Applications

Nickel-based alloys are required for wet corrosive applications if the conditions are so severe that stainless steels would corrode too fast or nonmetallic materials would have insufficient lifetimes due to erosion or other damage. Wet corrosive applications of nickel-based alloys include processing and handling of aggressive media in the chemical process industry, in environmental technology (e.g., flue gas desulfurization), and in marine and offshore technology. In a broader sense the use of unalloyed nickel as a coinage material also belongs to the category of wet

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7 Nickel-Based Alloys

corrosive applications. Unlike mechanical characteristics, corrosion resistance is not a materials property but a mode of behavior, i.e., it is always dependent on an interaction between the material and the surrounding corrosive medium. In addition, the kind and rate of this interaction may be governed by a third partner forming in between the two as a passivating or inhibiting surface film on the alloy. Normally the fundamentals of materials science and technology, as described in this Series, do not allow any predictions to be made with respect to the behavior of nickel-based alloys in wet corrosive applications. Even the results of standardized corrosion tests such as those on intercrystalline, pitting, and crevice corrosion only reflect the behavior of the alloy under the test conditions and are definitely not suitable for consideration as data on intrinsic materials properties. But empirical knowledge does allow us to conclude under which corrosive conditions nickel-based alloys will be required. In many cases, an additional direct test in the envisaged corrosive medium is indispensable for the ultimate decision as to which alloy should be applied and which should not. After corrosion resistance, with only a few exceptions, weldability is the second most important characteristic of nickelbased alloys in this field of application. Like corrosion resistance, weldability is not a property but a mode of behavior, resulting from the interaction between the material and the welding process applied. The materials are required to be weldable when using one of the current welding processes, because otherwise the vessels, pipes, ductings, flanges, etc., cannot be manufactured and joined together.

7.1.3 High-Temperature Applications Nickel-based alloys can be used at high temperatures where steels are no longer sufficiently creep resistant, i.e., at temperatures from above about 550 °C up to about 950 °C where long-time service is required and the 100 000 h creep-rupture strength is a criterion, or even up to about 1200 °C when shorter life times of, e.g., 1000 h or several years can be regarded as sufficient, as in the case of special industrial furnaces. Along with creep resistance, sufficient corrosion resistance to hot gaseous media is also an important requirement in this application. In most cases we are concerned with oxidation. In general, nickel-based alloys are superior to heat-resistant steels in cases where high resistance to cyclic oxidation is required. In various industrial hightemperature applications and in environmental technology, resistance to carburization, sulfidation, nitridation, and chlorination are also important materials characteristics. Apart from creep resistance, the requirement of special mechanical properties is not a primary reason for the use of nickelbased alloys. But once the mechanical properties have been determined, they have to be closely controlled in order to allow the estimation of maximum allowable stresses or to enable forming, stamping, and cutting operations under constant machining conditions. This is also true for wet corrosive (Sect. 7.1.2) and electrotechnical applications (Sect. 7.1.4). 7.1.4 Electro technical Applications Nickel-iron alloys with and without the addition of some other alloying constituents are finding applications as core materials for inductive components where extremely low energy losses are required and which cannot be met by means of

7.2 Composition, Structure, Properties, and Behavior

cheaper materials, e.g., iron-silicon alloys or nonmetallic ferrites. Magnetic permeability is the paramount materials property, being determined, for a given alloy composition, to a large extent by the grain size obtained in the final annealing, which is again a function of the density and distribution of nonmetallic inclusions caused by the metallurgical practice, i.e., deoxidation on melting of those alloys. Another electrotechnical application where nickel-based alloys are involved is in heating resistance elements. Here the specific electrical resistance and its dependency on temperature are important materials characteristics, but the resistance to oxidizing or other atmospheres under cyclic thermal conditions will determine the lifetime of the elements. In this application nickel-based alloys are preferred to competing iron-based materials if one or more of their characteristics, i.e., easier handling in manufacturing, better thermal stability, higher creep resistance, or better compatibility with isolating MgO powder and nitrogen furnace atmosphere, is a criterion. Spark plug electrodes are used for a combination of electrotechnical, high-temperature and, in most cases, automotive applications of nickel-based alloys. The dominant technical requirements are sufficient resistance to high-temperature gaseous corrosion under motor conditions along with high thermal conductivity and a combination of softness and suitable surface condition in order to allow a high tool life during electrode manufactures.

351

7.2 Composition, Structure, Properties, and Behavior 7.2.1 Nickel-Based Alloys for Wet Corrosive Applications 7.2.1.1 Composition, Structure, and Resulting Mechanical Properties

Main Alloying Constituents Table 7-1 lists the nominal chemical compositions of some typical nickel-based alloys for wet corrosive applications. First there is unalloyed nickel in its high carbon (max. 0.05 mass%) and special low carbon (LC) (max. 0.02mass%) forms. The nickel-copper alloy 400 and its age-hardenable version alloy k-500 follow and are widely used, whereas the use of the quasi-binary nickel molybdenum alloys B-2 and B-4 is limited to very specialized applications only. The nickel-chromium-iron alloys can be considered to be an extension of the austenitic stainless steels up to higher nickel contents. The same is true for the family of nickel - chromium - iron - molybdenum copper alloys, which contain molybdenum and copper as additional components besides chromium and iron as their major alloying constituents, whereas the family of nickel-chromium-molybdenum alloys has only a limited iron content. As can be seen from Table 7-1, the phase diagrams of primary importance in the case of nickel-based alloys for wet corrosive applications are the binary diagrams (Massalski, 1991) nickel-copper and nickel-molybdenum, and the ternary diagrams nickel-iron-chromium (Rivlin and Raynor, 1980) and nickel-chromiummolybdenum (Raghavan et al., 1984). Generally, the wrought nickel-based alloys for wet corrosive applications are designed to be single-phase austenitic, i.e., face-cen-

352

7 Nickel-Based Alloys

Table 7-1. Nominal chemical compositions of nickel-based alloys for wet corrosive applications. Main alloying constituents (mass%)

Designation Alloy

Trade name

Ni

Cr

Nickel 200 201

Nickel 99.2 LC-Ni 99.2

>99 >99

Nicorros Nicorros Al

64 65

Mo

Cu

Fe

Other

Nickel-copper 400 k-500

32 30

1.8 1

1 Mn 0.6 Mn 2.8 Al 0.45 Ti

1.8 3

0.7 Cr 1.3 Cr

Nickel- molybdenum Nimofer B-2 Nimofer B-4

69 66

Nickel - chromium - iron, Nicrofer 800 L 3220 LC

32

21

45

600 L 690

74 61

16 29

9 9

7216 LC 6030

28 27

Nickel - chromium - iron - molybdenum - copper,Nicrofer 28 31 3127 LC 27 31 3127 hMo 31 27

3.5 6.5

1.3 1.2

35 31

20 825 G-3 48

2.4 3.2 7 6.5

3.4 2.2 2.0 1.9

34 31 19 18

3620 Nb 4221 4823 hMo 4823 Ti

38 40 48 49

Nickel - chromium - molybdenum, NicroferhMo 625 6020 hMo 62 C-276 5716 hMoW 57 C-4 6616 hMo 66 22 5621 hMoW 59 59 5923 hMo 59

20 23 23 21

22 16 16 21 23

tered cubic during service, and the soft- or solution-annealed recrystallized condition is used in most cases. Due to the simplicity of the binary phase diagram of nickel-copper, this poses no serious problems in the case of the nickel-copper alloy 400. In contrast, for the nickel-molybdenum alloy B-2, according to Table 7-1, the intermetallic phase Ni 4 Mo forms below 840 °C so the alloy has to be maintained in face-

9 16 16 14 16

3 6 1 2 1

0.5 Ti 0.25 Al 0.25 Ti 0.25 Ti

1.7 Mn 0.2 N 0.2 Nb 0.8 Ti 0.3 Nb 1.9 Ti 0.6 Al 3.4 3.5 0.3 3

Nb W Ti W

centered cubic solid solution by quenching or rapid cooling from higher temperatures. This may not be so easy in practice in cases where heavy sections of plates or forgings have to be fabricated and where embrittlement, due to intermetallic phase formation, may cause cracking. Here the addition of some iron and chromium to the alloys, as indicated in Table 7-1, is very helpful since these elements are reported to

7.2 Composition, Structure, Properties, and Behavior

delay the intermetallic phase formation (Buth and Kohler, 1992). On the basis of this knowledge, alloy B-4 has been developed to offer easier fabrication and manufacture, and a reduced risk of stress-corrosion cracking during service, due to its increased iron and chromium content (Kohler and Heubner, 1994). The requirement of a face-centered cubic microstructure does not conflict with the corresponding ternary phase diagram in the case of the nickel-chromium-iron alloys reported in Table 7-1, but the situation is somewhat different from most of the nickel-chromium-iron-molybdenumcopper alloys and for the nickel-chromium-molybdenum alloys. Due to their composition they are close to the boundary which separates the austenite from multiphase regions where intermetallic phases will appear. Figure 7-1 shows the 1250 and 850 °C sections of the nickel-chromiummolybdenum phase diagram (Raghavan etal., 1984). Beyond the single-phase y austenite field, with increasing chromium content, the intermetallic phases designated 5, \x, P, and a exist. The \i phase forms

353

only at lower temperatures. These closepacked, hard and brittle phases separate the face-centered austenite from the bodycentered cubic chromium-molybdenum region. They reduce the ductility and take up the elements chromium and molybdenum, which are added to increase the corrosion resistance of the austenitic alloy matrix. The latter is especially effective if these phases precipitate from the solid austenite at temperatures below about 1000 °C, causing a depletion of chromium and molybdenum in the surrounding austenitic matrix. However, it has to be mentioned that if these intermetallics exist in a fine distribution in a limited quantity in a cast microstructure of the same overall composition, e.g., in weldments, their effect is much less pronounced and can be tolerated completely for practical applications (Heubner etal., 1990; Kirchheiner etal., 1992). In order to avoid the multiphase fields of Fig. 7-1, the influence of the additional alloying elements on the position of the phase boundary has to be evaluated. This can be done approximately by means of

Mo

Mo

20

Ni

40

20 Mass %

60

Ni

20

40

60

Mass %

Figure 7-1. Isothermal sections of the N i - C r - M o ternary system at (a) 1250°C and (b) 850°C, according to Raghavan et al. (1984).

354

7 Nickel-Based Alloys

phase computation attempts. A more recent approach (Morinaga etal., 1984), consists of computing the average energy level of the atomic d-orbital of the alloy according to

(Md)i

M, =

(7-1)

where Xt is the molar fraction of the component / (there are n components) and Md the average energy level of the d-orbital. According to Morinaga et al. (1984), precipitation of the a-phase will take place if Md, as a function of temperature, exceeds the values resulting from = (6.25 x

(7_2)

0.834

where T is the temperature in Kelvin. As Fig. 7-2 demonstrates, within the 850 °C section of the nickel-chromiummolybdenum system there is good correlation of this computation with the experimentally established y-phase boundary.

80

40

60

20

80

40

60 Mass %

Figure 7-2. Isothermal section of the N i - C r - M o ternary system at 850 °C. Phase equilibria according to experimental results (Raghaven et al. 1984), and a trace of the iso-Md line for M d = 0.905 (850 °C) ( ) according to Eq. (7-2).

60 50 Ac 40

^30

Experimental * • Calculated Ni-Cr-Mo line Calculated with 3 . 5 % Nb

10

¥ 20

30 40 Mass % Cr

50

y—— 60

Figure 7-3. Isothermal section of the N i - C r - M o ternary system at 1250 °C. Phase equilibria according to experimental results (Raghaven et al., 1984), ineluding trace of the iso-Md line for Md = 0.929 (1250 °C) according to Eq. (7-2) and displacement of this line as a result of an addition of 3.5% niobium (Kohler and Heubner, 1992).

According to Fig. 7-3 greater differences result for the 1250 °C section of the ternary nickel-chromium-molybdenum phase diagram. Figure 7-3 shows at the same time how, according to the phase computation, the y-phase boundary will be shifted if 3.5% niobium is added to nickel-chromium-molybdenum alloys (Heubner and Kohler, 1992). The nickel-chromium-molybdenum alloys and most of the nickel-chromiumiron-molybdenum-copper alloys are designed in such a way that the desired austenitic microstructure can only be established at high temperatures, so that the alloys have to get a final high-temperature annealing treatment followed by rapid cooling in order to suppress precipitation of the intermetallics. This is reflected in Table 7-2, which indicates typical final annealing treatments for the alloys mentioned in Table 7-1. Since wet corrosive applications are limited by temperature, the intermetallic phases are not usually ex-

355

7.2 Composition, Structure, Properties, and Behavior

Table 7-2. Heat-treatment guidelines for soft- and solution-annealing of nickel-based alloys for wet corrosive applicationsa. Designation

Soft-annealing b

Solution-annealing 0

Temp.b (°C)

Cooling0

air air

-

-

700-900 -

air -

980-1040

water

-

-

1050-1080

water air jet

water water/air water

-

-

Nickel - chromium - iron - molybdenum - copper, Nicrofer 28 3127 LC 3127 hMo 31 3620 Nb 20 980-1010 825 4221 920-980 G-3 4823 hMo 4823 Ti 48 d -

water water/air -

1080-1150 1130-1180 1130-1170 1150-1190

water water water water

Nickel - chromium - molybdenum, Nicrofer hMo 950-1050 6020 hMo 625 5716 hMoW C-276

air air

1100-1180 1080-1135

-

1050-1080

-

1105-1135

-

1100-1180

air water air jet water air jet water air jet water air jet

Alloy

Trade name

Temp. (°C)

Nickel 200 201

Nickel 99.2 LC-Ni 99.2

700-850 700-850

Nickel-copper 400 Nicorros k-500c Nicorros Al Nickel-molybdenum B-2 Nimofer Nickel- chromium - ion, Nicrofer 800 L 3220 LC 600 L 7216 LC 6030 690

C-4

6616 hMo

22

5621 hMoW

59

5923 hMo

980-1050 920-980 950-1050

-

-

Cooling

a Heubner (1987b), data from Krupp VDM GmbH; b a suitable holding time at temperature is usually 3-4 min/ mm; c in bright-annealing of strip the protective atmosphere provides a sufficient amount of cooling in most cases; d solution-annealing followed by subsequent age-hardening treatment, according to Table 7-3.

pected to precipitate during the alloys9 service. They may, however, precipitate during manufacturing operations when vessels, ductings, and other equipment have to be made from plates, sheets, and tubes by means of welding and hot forming, where the materials are partially or wholly

heated to intermediate temperatures and cannot receive a subsequent high-temperature final anneal due to their size or for other reasons. It is, therefore, essential to have sufficient knowledge about the time-temperature-precipitation (TTP) behavior and the

356

7 Nickel-Based Alloys

Start of precipitation 1100- at qrainboundaries

©

1000o

Z3

ro t_ QJ d,

E QJ

t—

/© 900- ©

V\\© \ 700\

800-

;te€)

© ©

\© ^©

©

6>s

© © © ^ ©

oo oo G G

-£

600 500 0.1

© N> © ©

t

^^ 0^)

0.3

1

3 10 Time (h)

Within grains

OO

oo 00 oo oo Q

o © ©

100

Figure 7-4. Isothermal time - temperature - precipitation (TTP) diagram for alloy 59, established by means of optical microscopy for a melt with (mass%) 60.8 Ni, 22.7 Cr, 15.6 Mo, 0.3 Fe, 0.03 Si, and 0.005 C (Agarwal et al., 1991). Key: O free from precipitates, @ precipitation starting at grain boundaries, © some grain-boundary precipitation, © grain-boundary precipitation still incomplete, O precipitation covers boundaries completely, £> some precipitation within the grains (max. 1 % of grain volume), Q precipitation covering 1-5% of grain volume, # precipitation covering more than 5%. © NACE, reprinted with permission.

resultant change in the characteristics of the material. As an example, Fig. 7-4 shows the time-temperature-precipitation behavior of the nickel-chromiummolybdenum alloy 59 (Agarwal et al. 1991; Heubner and Kohler, 1992). As might be expected, precipitation starts at the grain boundaries and the first precipitates will be observed under the optical microscope after about 6 min at 900 °C. Subsequent precipitation takes place in the temperature range between about 550 and 1050 °C. The precipitating phase in this alloy is composed of (mass%) 31.8 Mo, 20.3 Cr, and 48.2 Ni and is therefore (in accordance with the nickelchromium-molybdenum phase diagram) interpreted as j^i-phase. Although the carbon content of this alloy is about 0.005 mass%, an accompanying precipitation of carbides, e.g., M 6 C, is likely (Heub-

ner and Kohler, 1992). The ductility in terms of notch impact strength in the solution-annealed condition is very high. Values of up to 300 J have been measured on ISO V-notch samples at ambient and at cryogenic temperatures ( —196°C). As Fig. 7-5 demonstrates, these high values are reduced by precipitation of the intermetallic phase on annealing at intermediate temperatures, but they are well above 100 J even after 8 h annealing. The time-temperature-sensitization (TTS) diagram shown in Fig. 7-6 indicates an intergranular penetration of more than 0.05 mm (2 mils) in the ferric sulfate sulfuric acid test according to ASTM G 28 A (24 h) only after about 2 h at 850 °C. This time span is longer than that observed on any other N i - C r - M o alloy so far (Kirchheiner et al., 1990). As another example, Fig. 7-7 shows the time-temperature-precipitation behavior of alloy G-3 (Heubner and Kohler, 1992). Again precipitation starts at the grain boundaries, but two different reactions occur. One is the precipitation of a-phase, discernible under the optical microscope after about 20 min at 900 °C, and the other

1150 11001000 950 900 850 800 700 Temperature (°C)

Figure 7-5. Isothermal time - temperature - notch impact strength diagram for alloy 59, established on ISO V-notch samples at ambient temperature for a melt with (mass%) 60.8 Ni, 22.7 Cr, 15.6 Mo, 0.3 Fe, 0.03 Si, and 0.005 C (Agarwal et al., 1991). © NACE, reprinted with permission.

7.2 Composition, Structure, Properties, and Behavior 1100

357

the ASTM G-28 A standard test (Heubner and Kohler, 1992). Accompanying Elements: Carbon and Others

10 Time (hfc

100

Figure 7-6. Isothermal time -temperature -sensitization diagram for alloy 59, established on a melt with (mass%) 60.8 Ni, 22.7 Cr, 15.6 Mo, 0.3 Fe, 0.03 Si, and 0.005 C after 24 h boiling in 42 g/1 ferric sulfate50% sulfuric acid (ASTM G-28, method A), indicating the field where intergranular penetration (IP) exceeds 0.05 mm (2 mils) (Agarwal et al. 1991; © NACE, reprinted with permission).

~T

0.3 0.5 1

Start of precipitation T~at grain boundaries n

3 10 Time (h)

100

Figure 7-7. Isothermal time - temperature - precipitation behavior of alloy G-3, established by means of optical microscopy for a melt with (mass%) 47.6 Ni, 22.6 Cr, 19.7 Fe, 7.1 Mo, 2.0 Cu, 0.4 Mn, 0.3 Nb, 0.03 Si, and 0.01 C (Heubner and Kohler, 1992). For the symbols see legend to Fig. 7-4.

is carbide precipitation, becoming apparent under the optical microscope after about 6 min at 650 °C. However, the kinetics of this carbide precipitation are so slow that a much longer time, about 3 h, is needed for intercrystalline corrosion with more than 50 |im penetration according to

In continuation of the preceding comments on Fig. 7-7 which should be read in this context, it has to be mentioned that precipitation of carbon in the form of graphite may occur in unalloyed nickel (alloy 200 in Table 7-1), and that precipitation of carbon in the form of carbides may occur in nickel-chromium-iron-molybdenum-copper and in nickel-chromiummolybdenum alloys in addition to precipitation of the intermetallic phases. Graphite precipitation at intermediate temperatures (about 300 °C and above) in unalloyed nickel can be a nuisance since it takes place preferentially at grain boundaries and thus impairs the ductility. This can be prevented by using the low carbon or LC-grade alloy 201, which is why this grade was created. Apart from nickel-copper alloys, the nickel-based alloys for wet corrosive applications, as mentioned in Table 7-1, are usually manufactured low in carbon nowadays, with carbon contents of 0.025 mass% maximum or even much less, a carbon content of about 0.005 mass% being typical for the nickel-chromiummolybdenum alloys C-276, C-4, and 59. This is due to the progress made in melting in the 1960s with the introduction of argon oxygen decarburization (AOD) and vacuum oxygen decarburization (VOD). The alloys created before this time were originally designed to have more carbon. Typical examples are alloys 825 and 20 in Table 7-1. In alloy 825 the high titanium content was introduced to absorb the carbon by forming TiC in a final, so-called stabilizing soft-anneal at about 950 °C (see Table 7-2). Generally, the temperature used for these

358

7 Nickel-Based Alloys

final stabilizing annealing treatments is high enough to allow complete recrystallization but, at the same time, as low as possible in order to precipitate most of the carbon. Due to the forces of tradition in the life of materials' specifications, this alloy continues to be manufactured with a high titanium content and the final stabilizing anneal, even if its carbon content is now, in most cases, limited to 0.015 mass% maximum. However, on a special request for higher strength, which might be wanted occasionally for offshore applications, a high carbon version with 0.04-0.05 mass% carbon can still be manufactured without any other change being necessary. Also alloy 20, in its version as Nicrofer 3620 Nb, is now manufactured with a maximum carbon content of 0.015 mass%. In this case niobium is the so-called carbon stabilizer, and it was possible to lower the niobium content along with the carbon. Niobium is undesirable as a carbon stabilizer, since niobium carbides have a strong tendency to form on the grain boundaries in nickelbased alloys in such a way that the ductility and resistance to intergranular corrosion are sharply reduced. The most prominent surviving and still going strong example of a niobium-stabilized nickel-based alloy is alloy 625 for wet corrosive applications, so care must be taken during manufacture that most of the niobium carbide is precipitated before final recrystallization takes place. Other accompanying elements which have to be closely controlled are sulfur, phosphorus, and silicon. Sulfur in excess of about 0.005% impairs the resistance to pitting corrosion and silicon promotes precipitation of the aforementioned intermetallic phases. Typical silicon contents of today's nickel-chromium-molybdenum alloys are therefore around 0.04mass% (Agarwal etal., 1991).

Age-Hardening Alloy Constituents While the alloys are used in a soft- or solution-annealed condition in most cases, there are some applications where use is made of cold-worked, deformation-hardened material, e.g., in the oil and gas business where a major requirement is for a combination of suitable corrosion resistance and high strength/weight ratio for downhole applications. In these cases agehardenable alloys are also required. Age hardening is achieved by the precipitation of a y'(Ni,Cu)3Al phase, as in the case of alloy k-500, (Volk, 1970) or a y/Ni3(TiAl) phase, as in the case of alloy 48 (Bryant et al, 1990). Figure 7-8 shows the y/y' mi-

Figure 7-8. y/y' microstructure of (a) alloy k-500 and (b) alloy 48 in TEM bright and dark field illumination, respectively.

7.2 Composition, Structure, Properties, and Behavior

crostructures found in the fully age-hardened condition for both types of alloy. For these cases, the quaternary N i - C r Al-Ti system (Taylor, 1956), as shown in Fig. 7-9, has to be looked at. However, due to the simultaneous precipitation of additional phases like carbides and a, the situation might be much more complex in practice (Bryant et al., 1990). For example, in the case of alloy 48 the aging treatment has to be conducted at sufficiently low temperatures in order to avoid the temperature range in which embrittling high-temperature grain-boundary precipitation takes place. On the other hand, the aging temperature has to be sufficiently high so that the maximum strength can be achieved. Consequently, the aging treatment has to follow a rather narrow route. Typical thermal treatments for the two age-hardenable alloys in Table 7-1 are given in Table 7-3.

359

Table 7-3. Some guidelines for age-hardening treatments subsequent to solution-annealing or hot forming applicable to nickel-based alloys for wet corrosive applicationsa. Designation

Age-hardening treatment

Alloy

Trade name

Temperature (c

k-500

Nicorros Al

1)1 595 16 II 540 6 III 480 8 3-step anneal with furnace cooling between I and II, and II and III 2)1 595 16 II Cooling to 480 °C at 8-14°C/h 3)1 640 2 II Furnace cooling over 10 10hto480°C

48

Nicrofer 4823 Ti

1)1 750 II Cooling to 600 °C at 15°C/h III air cooling 2)1 730 II Cooling to 580 °C at 15°C/h III Air cooling 3)1 730 II Cooling to 580 °C at 15°C/h III Air cooling

Time (h)

4

8

4

Heubner (1987 b), data from Krupp VDM GmbH.

NiAl

Ni 3 Al

Figure 7-9. The nickel - chromium - aluminium - titanium phase diagram at 750 °C (Taylor, 1956).

The widely used high-temperature superalloy alloy 718 (see Sect. 7.2.2.1) is also used for wet corrosive applications. Due to its high niobium content alloy 625 also has a certain capability for being age hardened (Kohler, 1991), and can be improved further by additions of titanium (Frank and DeBold, 1988; Hibner, 1991). In addition,

360

7 Nickel-Based Alloys

an age-hardenable version of alloy 825 exists designated alloy 925 (Ganesan et al., 1987).

Table 7-4. Mechanical properties (minimum values) of nickel-based alloys for wet corrosive applications at ambient temperaturesa. Designation

P0.2

A5*

(N/mm2)

(N/mm2)

/o/ \ v /o)

100 80

370 340

40 40

Nickel-copper 400 Nicorros k-500 Nicorros Ale

180 690

450 965

35 20

Nickel- molybdenum B-2 Nimofer

350

760

40

450 550 600

35 30 45

Mechanical Properties

Alloy

Trade name

Although the corrosion behavior and, in most cases, the response to welding are the decisive characteristics of nickel-based alloys for wet corrosive applications, their mechanical properties have to be known in order to assess their use as engineering materials. Table 7-4 provides an initial guideline of the mechanical properties that might be expected in tensile testing of nickel-based alloys at ambient temperatures. The minimum values which can be achieved, even with coarse-grained material, are indicated. Indeed, with the exception of only a few wet corrosive applications of deformation-strengthened or agehardened alloys, nickel-based materials are used for wet corrosive applications in their soft- or solution-annealed, i.e., recrystallized condition. So the mechanical properties will be determined initially by the solution hardening and by the grain size. The degree of solution hardening is fixed by the chemical composition with a small band of scatter according to the users' specifications. The grain size may vary between about 50 \xm and 250 |im, according to the preceding manufacturing treatment and the final solution-anneal required. In many cases, the grain size of nickel-based alloys for wet corrosive applications will be 125-180 |im. In general, the grain size can be controlled more closely if thinner sections are to be manufactured, whereas in the case of heavy plates or forgings large grains may be difficult to avoid. However, during the manufacture of thinner sections, e.g., strip or thin sheet, a flattening or straightening operation might be the last manufacturing step, which gives

Nickel 200 201

Nickel 99.2 LC-Ni 99.2

Nickel-chromium-iron, Nicrofer 800 L 3220 LC 180 600 L 7216 LC 180 690 6030 300

Nickel- chromium - iron - molybdenum- copper, Nicrofer 28 3127 LC 220 500 35 31 3127 hMo 320 690 50 20 3620 Nb 240 550 30 4221 825 240 550 30 G-3 4823 hMo 240 620 45 48 4823 Tif 800 1100 20 Nickel-chromium-molybdenum, Nicrofer hMo 6020 hMo 625 415 830 C-276 5716 hMoW 310 750 C-4 6616 hMo 305 700 22 5621 hMoW 310 690 59 5923 hMo 350 710

35 30 40 45 45

a

Krupp VDM GmbH (1992);b yield strength;c rupture strength; d tensile elongation; e hot-finished, age-hardened; f solution-annealed, age-hardened.

these materials a slight additional deformation hardening, which does not occur with thick sections. Consequently the typical yield strength values will be well above those indicated in Table 7-4 as a minimum, and thin sections can more easily be provided with higher yield strengths than thick ones. With exceptions, the ductility is not generally a problem with nickel-based alloys

7.2 Composition, Structure, Properties, and Behavior

for wet corrosive applications, but it may be seriously impaired by grain-boundary precipitation of graphite, carbides, or intermetallic phases. This is especially true for the age-hardenable alloys, but also for unalloyed nickel, so that the low-carbon nickel grade has to be used if the nickel is exposed to temperatures of 300 °C and above during service. Also the maximum temperatures of operation, which can be allowed according to the safety requirements, for the long-time service of nickelbased alloys in wet corrosive pressure vessel applications are determined by their tendency to internal grain-boundary precipitation (see Figs. 7-4 and 7-7) and the accompanying loss of ductility. According to German TUV pressure vessel regulations (TUV is the authority responsible for safety checks), the maximum service temperatures are 400 °C (alloys B-2, 800 L, C-4, and C-276), 425 °C (alloy 400), 450 °C (alloys 600 L, 625, and 825), and 600 °C (alloy 201). Finally, it has to be mentioned that agehardenable alloys, e.g., alloy k-500, may experience a loss of ductility if, during seawater service, an uptake of hydrogen takes place due to direct galvanic coupling with a less noble material, e.g., carbon steel, or due to an inadequate cathodic protection system (Bryant and Wallis, 1988). The same might happen if sulfur-containing lubricants have been used for direct couplings with carbon steel in downhole equipment of oil and gas wells and those lubricants start to decompose when the temperature increases. One special form of ductility loss, stress-corrosion cracking, has to be considered as well in this context and will be addressed in the following section.

361

7.2.1.2 Composition, Corrosion Behavior, and Resulting Special Uses

Some more or less exhaustive books (Volk, 1970; Friend, 1980; Heubner, 1987 a, 1995) and monographs (Klarstrom etal., 1987; Brill, 1992) on the corrosion behavior of nickel-based alloys exist, so this section concentrates on the more important aspects. Special uses are indicated in actual brochures (e.g., Krupp VDM GmbH, 1992). In the following the terms oxidizing and reducing solutions or acids will be used. Although not scientifically correct these terms are broadly used today. Oxidizing acids are those where anions like NO^ or the oxygen in aqueous solution constitute the oxidizing agents, whereas in so-called reducing acids only the H + ion acts as an oxidizing agent. There are many mixed situations, e.g., reducing acids with an addition of dissolved oxygen or oxidizing ions. Nickel Unalloyed nickel in its common commercial forms for wet corrosive applications is mostly nickel alloy 200 and the low-carbon grade nickel 201, as mentioned in Table 7-1. The purity is at least 99 mass%, with iron and manganese making up most of the balance. Higher purity grades also exist. As Friend (1980) points out, the standard reduction potential of nickel is more noble than that of iron and less noble than that of copper. Because of nickel's marked overpotential for hydrogen evolution, it does not discharge hydrogen readily from any of the common nonoxidizing acids and a supply of oxygen is usually necessary for significant corrosion to proceed. However, nickel has the ability to protect itself by the development of a passive oxide film, and in

362

7 Nickel-Based Alloys

service conditions it is often observed to be more noble than copper. The most outstanding feature of unalloyed nickel is its high resistance to alkalies such as sodium and potassium hydroxide at high concentrations and temperatures, and to nonoxidizing alkaline and neutral salts. Whereas low-alloy austenitic steels can be used at NaOH concentrations of less than 50% and temperatures below 100°C, unalloyed nickel is the material most worthy of consideration if more aggressive conditions prevail. As explained above, at temperatures higher than 300 °C, the low-carbon grade alloy 201 has to be used. In oxygen-free alkaline media, the resistance to so-called caustic cracking also increases with the nickel content of the alloys, unalloyed nickel being fully resistant to this type of corrosion. However, in oxygen-bearing alkaline media, both the nickel and the chromium contents must be high, i.e., nickel-chromium-iron alloys have to be used in order to achieve adequate resistance to caustic cracking. While nickel can be rapidly corroded by oxidizing acids, it is sufficiently resistant to most nonaerated organic acids and other organic compounds. On the other hand, nickel has a high degree of resistance to different kinds of water including rapidly flowing seawater. Therefore typical applications also include specific components for plastic and food production. Nickel- Copper Alloy 400 and its age-hardenable version alloy k-500, as mentioned in Table 7-1, are the most commonly used nickelcopper alloys. Their most important characteristic for today's technology is their high resistance to the corrosive attack of various waters. So the main applications of alloy 400 include feedwater heaters, steam

generators, and other heat exchangers in energy producing plants, splash zone sheathing of risers, and platform steel legs in offshore oil and gas production. Alloy k-500 is also mainly used for seawater and offshore technology where high strength is required in addition to corrosion resistance, including applications such as pump and propeller shafts, valves and bushes, instrumentation casings, as well as other accessories in oil and gas production. Additionally, due to its high nickel content, alloy 400 is resistant to most alkalies except ammonium hydroxide, to most nonoxidizing alkaline and neutral salts, and to halogen gases at elevated temperatures. Due to its resistance to dilute hydrochloric acid, alloy 400 can withstand the attack of hydrolized, nonoxidizing chloride salts. Therefore other applications are in plants for the concentration and crystallization of salt and for the production of vinyl chloride monomer (VCM), and for overhead crude condensers in oil refineries. Nickel- Molybdenum Molybdenum metal has very good resistance to corrosion by nonoxidizing solutions of hydrochloric, phosphoric, and hydrofluoric acid at most concentrations and temperatures, and to boiling sulfuric acid up to about 60 % concentration. A molybdenum addition to nickel and nickel alloys improves the resistance to all of these environments approximately in proportion to the amount of molybdenum added. The nickel-molybdenum alloys, normally containing 26-35 mass% molybdenum, are among the few metallic materials that are resistant to corrosion by hydrochloric acid at all concentrations and temperatures (Friend, 1980). In today's technology the nickel-molybdenum alloy B-2 given in Table 7-1 will

7.2 Composition, Structure, Properties, and Behavior

363

Corrosion rates in parentheses are in mm/year 'Over 50 mpy (over 1.27) lOmpy '(1.27) 20 mpy ^ S , ^ (0.51) s . 20-50 mpy

-200

-50 0

10

20

30

40 50 60 70 Mass% H2S04

80

Figure 7-10. Isocorrosion diagram of alloy B-2 in quiet unaerated sulfuric acid solutions (Lee and Hodge, 1976; © NACE, reprinted with permission).

90 100

find application in most cases. It is particularly used for equipment to handle hydrochloric acid at all concentrations and temperatures, including at its boiling point. Other applications are in hydrogen chloride gas and in acetic, sulfuric, and phosphoric acids. With respect to sulfuric acid, the isocorrosion diagram shown in Fig. 7-10 exemplifies the broad usefulness of this alloy. Due to its low carbon and silicon contents, it exhibits good resistance to knife-line and heat-affected zone attack, and can therefore be used in the as-welded condition in most cases. Alloy B-4 is a more advanced version of alloy B-2, offering more ease in fabrication and reduced risk of stress-corrosion cracking during service (Kohler and Heubner, 1994). Nickel- Chromium-Iron Chromium is added to nickel and nickel alloys to improve the resistance to strongly oxidizing environments such as nitric and chromic acids, and to some acid solutions containing oxidizing salts, except those involving chlorides. The resistance is roughly proportional to the amount of chromium added, so long as it is above about 10 mass%. In the case of aerated acid solutions, particularly at higher temperatures,

chromium additions of over 18 mass% are to be preferred (Friend, 1980). Since a more universal resistance to wet corrosion, especially to chloride-bearing media, requires the presence of additional alloying elements like molybdenum, the nickel-chromium-iron alloys indicated in Table 7-1 find usage preferentially in special applications. Alloy 800L is the low-carbon variant of the important high-temperature alloy 800. Due to its excellent resistance to corrosive attack by high-purity, high-temperature water, it has become the standard material for steam generators and feedwater heaters in nuclear and fossil-fired power plants. For nuclear applications a special version is available with a restricted cobalt content of 0.01 mass% maximum, but the alloy is also used in the chemical industry, e.g., ammonia plants. Alloy 600L is also a low-carbon version of a widely used, high-temperature material, alloy 600. This alloy is resistant to ammonium hydroxide, amines, and hot fatty acids. It is used in food processing and pharmaceutical applications where a high degree of corrosion resistance as well as complete freedom from chloride stresscorrosion cracking are required. As with alloy 800L, an important application of

364

7 Nickel-Based Alloys

alloy 600L is for steam generator tubing in nuclear power plants of the pressurizedwater reactor (PWR) type. Alloy 690 was developed in the search for an improved version of alloy 600L for handling pressurized hot-water, i.e., steam generator tubing in nuclear power plants of the PWR type, where boiling and steam blanketing at heat transfer surfaces can give rise to very high local concentrations of caustic, and as a result caustic-induced stress-corrosion cracking or caustic cracking may occur. It is also a useful construction material in cases where aerated alkaline media are to be handled and where high levels of both nickel and chromium are necessary for sufficient resistance to caustic cracking (Sedriks, 1979). Due to its high chromium content, alloy 690 also has high corrosion resistance to the heavily oxidizing media that may be produced in nuclear fuel reprocessing in which waste is dissolved in nitric acid, especially when it is contaminated by fluorides (Heubner and Kirchheiner, 1987). Nickel- Chromium - Iron - Molybdenum Copper Copper additions to stainless steels and nickel-chromium-iron-molybdenum alloys have been shown to increase their resistance to reducing acids. By experience the copper content is restricted to 4 % max. Microprobe analysis revealed a copper layer being formed on the surface during corrosive attack as the reason for this improved corrosion resistance. This copper precipitate may act by facilitating the cathodic reaction so that the alloy can be passivated more easily. Additionally this copper layer may act as a diffusion barrier, thus lowering the metal dissolution rate (Heubner et al. 1985). Too high amounts of copper are suspected of causing a de-

crease in the resistance to localized corrosion (Heubner and Rockel, 1986; Renner et al., 1986) so that, with the exception of alloy 20, the copper contents of the nickelchromium-iron-molybdenum-copper alloys are limited to not much more than about 2%, as Table 7-1 shows. The other main alloying elements chromium and molybdenum have been mentioned in the foregoing sections. As in stainless steels, chromium is the paramount alloying constituent for providing passivity (Sedriks, 1979) whereas molybdenum is very essential for the alloys' resistance to pitting and crevice corrosion in chloridebearing media, being about three times as effective as chromium with respect to pitting and more than twice as effective as chromium with respect to crevice corrosion (Renner et al., 1986). Besides nickel, iron is another base element of those alloys. Increasing the nickel content by substituting nickel for iron increases the alloy's solubility for chromium and molybdenum, and so the alloys can be made more corrosion resistant. In addition, nickel has a positive influence of its own, so the corrosion resistance in reducing acids is improved with increasing nickel content (Sedriks, 1979). Furthermore, an increasing nickel content provides increasing resistance to chloride-induced stress-corrosion cracking (Rockel, 1987). The development of the nickel-chromium-iron-molybdenum-copper alloys mentioned in Table 7-1 was aimed especially at obtaining high resistance to corrosive attack by sulfuric and phosphoric acid solutions over a broad range of concentrations and temperatures, as much under oxidizing as under reducing conditions. The veteran in the group is undoubtedly alloy 825, which is still a successful material in today's technology. It is certainly the alloy's position in the central area of the

7.2 Composition, Structure, Properties, and Behavior

nickel-chromium-iron phase diagram and the only moderate dotation with additional alloying elements which make this alloy a multitalented workhorse in the chemical process industry and in offshore applications. It can also be used under more aggressive oxidizing conditions e.g., in nitric acid (Heubner and Kirchheiner, 1987), and is a useful material for handling caustic media. Alloy 20, similar to alloy 825, is outstanding because of its high copper content. Typical uses are heat exchangers and piping systems in sulfuric acid and phosphoric acid plants. Alloy G-3 and its age-hardenable version, alloy 48, differ from alloy 825 mainly insofar as the nickel and molybdenum contents are increased, providing both materials with improved resistance to pitting and crevice corrosion. One typical application of alloy G-3 is in flue gas desulfurization, whereas alloy 48 has been designed for use in sour oil and gas fields (sour gas is natural gas containing hydrogen sulfide). The greatest difference between alloy 28 and alloy 825 is the increased chromium content which provides alloy 28 with the special potential for being used under oxidizing conditions. Alloy 28 is, indeed, of interest where on the one hand resistance to nitric acid attack is required and on the other hand a high resistance to pitting and crevice corrosion has to be guaranteed, e.g., where deposits and incrustations may occur during service (Heubner and Kirchheiner, 1987). However, typical applications are in heat exchangers handling sulfuric and phosphoric acid contaminated by chlorides and fluorides (Kirchheiner et al., 1989). This alloy is also used in sour oil and gas-production tubing. Alloy 31 constitutes the most recent creation within the group of nickel -chromium-iron—molybdenum—copper alloys (Heubner et al., 1989 a, 1991 b). It exhibits

365

excellent resistance to pitting and crevice corrosion in neutral and acidic aqueous solutions, coming close to that of higher nickel alloyed materials. In seawater its pitting potential decreases to a greater extent only above 75 °C. Its resistance to attack by aggressive oxidizing media, such as boiling azeotropic nitric acid, is not impaired by its high molybdenum content. On the other hand, it exhibits high resistance to attack by reducing media, e.g., hot sulfuric acid, as shown in Table 7-5. The combination of high resistance to localized corrosion in hot chloride-bearing cooling waters, including seawater, with high resistance to corrosion by aggressive oxidizing and reducing hot corrosive media is a unique feature of alloy 31. Nickel- Chromium - Molybdenum As is to be expected from the corrosion behavior of nickel-chromium alloys in oxidizing media and of nickel-molybdenum alloys in reducing media, nickel-based al-

Table 7-5. Corrosion rate of alloy 31 in sulfuric acid (20-80%) immersion tests over 21 days 3 H 2 SO 4 Concentration (mass%)

Temperature (°C)

Corrosion rate (mm/year)

20

60 80 100

0.001 0.002 0.007

40

60 80 100

0.002 0.004 0.016

60

60 80 100

0.002 0.010 0.025

80

60 80 100

0.004 0.019 >6.0

Heubner et al. (1991 b).

366

7 Nickel-Based Alloys

loys rich in both chromium and molybdenum show a useful resistance to corrosion in both oxidizing and reducing environments. The first material of this type, alloy C, was introduced in the 1930s and is now obsolete. Alloy C-276 (see Table 7-1) is a modification of this alloy. The modification consisted basically of reducing both the carbon and the silicon content more than 10-fold, to very low levels of typically 0.005% and 0.04% respectively, in order to reduce the susceptibility to intergranular corrosion. To further improve upon the thermal stability of alloy C-276, a variant of this alloy was developed in the 1970s called alloy C-4 (Table 7-1). This development consisted of reducing the iron and omitting the tungsten in order to virtually eliminate any grain-boundary precipitate formation in the time periods likely to be encountered in most hot-working operations. Alloy 22 was developed later, and alloy 59 was introduced in 1990, showing improvements over both alloy C-276 and alloy 22 in certain environments (Agarwal etal., 1991). The tungsten content of alloys C-276 and 22 acts in a similar way to molybdenum with respect to the corrosion behavior, but the same amount of molybdenum (mass%) is a little less deleterious to the corrosion resistance in oxidizing media and much more effective than tungsten in reducing media and with respect to crevice corrosion (Heubner et al., 1989 c). The effect of niobium on the corrosion behavior of alloy 625 is less well documented, but it may be supposed that the niobium acts in a similar way to the molybdenum. Figure 7-11 shows the corrosion rate of four different nickel-chromium-molybdenum alloys in a boiling, strongly oxidizing sulfuric acid test solution. As is to be expected from the foregoing, the corrosion rate decreases with increasing chromium

X

Alloy C-276 (Ni57 Mo 16 W3.5 Fe6)

4 3-

X Alloy C-4 (Ni66 Mo 16 Fe3)

21 -

14

Alloy 22 (Ni 57 Mo 13 W 3 . 2 F e 3 ) X

16

18 20 Mass% Cr

AU (Ni

22

59

59 Mo 16 Fe 1)

24

Figure 7-11. Corrosion behavior of nickel-chromium -molybdenum alloys in a solution according to ASTM G-28, Method A: 50% H 2 SO 4 + 42g/l Fe2(SO4)3 • 9H 2 O (Kirchheiner et al, 1990; © NACE, reprinted with permission).

content of the alloys, i.e., alloy 59 behaves best under these test conditions. The different behaviors of alloys C-276 and C-4 is mainly due to the tungsten content of alloy C-276, which makes alloy C-276 less corrosion resistant under these oxidizing test conditions. On the other hand, alloy C-276 is superior to alloy C-4 in sulfuric acid condensates contaminated with chlorides, as may occur, e.g., in flue gas desulfurization plants. This is demonstrated by the data on the critical pitting temperature (cpt) and the critical crevice corrosion temperature (cct) in the so-called "Green Death" solution, as shown in Table 7-6. Under these conditions alloy C-276 behaves better than alloy C-4 and alloy 625, but alloy 59 outperforms them both (Kirchheiner et al., 1992). Alloy 59 also shows promising behavior in other reducing media. Figure 7-12 presents the isocorrosion diagram of alloy 59 in hydrochloric acid. Over the whole range (up to 40% HC1) the 0.13 mm/y (5 mpy) isocorrosion plot is at temperatures above 40 °C (104 °F), demonstrating that alloy 59 is, with respect to this behavior, a more versatile material than the other nickel-chromium-molybdenum alloys C-276, 22, and 625 (Agarwal et al., 1991).

7.2 Composition, Structure, Properties, and Behavior

367

Table 7-6. Critical pitting (cpt) and crevice corrosion temperatures (cct) in sulfuric acid solutions with chloride additions a. Alloy

Solution

"Green Death" + 7% H 2 SO 4 + 3 vol.% HC1 + 1% CuCl2 + 1% FeCl 3 -6H 2 O a

59

22

C-276

C-4

625

>120 110

120 110

115-120 105

100 85-95

100 85-95

cpt (°C) cct (°C)

Kirchheiner et al. (1990).

As a result of their high corrosion resistance, alloys C-276 and 59 find numerous applications, e.g., as components in critical chemical and petrochemical processes, in the pulp and paper industry, is downhole equipment for sour-gas production, and as components in flue gas desulfurization and related applications. Alloy C-276 proved to be resistant to chloride-induced stresscorrosion cracking in today's most agressive downhole sour gas media, even in the presence of elemental sulfur, but it is sensitive to hydrogen-induced stress-corrosion cracking in media of this kind below about 100 °C, which is attributed to the alloy's low iron content. Alloy C-4 has a similar corrosion behavior to alloy C-276, but

68 20 HCl J % )

Figure 7-12. Isocorrosion diagram of alloy 59 in quiet unaerated hydrochloric acid solutions (Agarwal et al., 1991; © NACE, reprinted with permission).

with somewhat less overall corrosion resistance. On the other hand, due to its other solidification characteristics and improved thermal stability, it may be easier to weld and to manufacture into various shapes. Alloy C-4 is a popular material with the chemical process industry in Germany, whereas alloy C-276 is more popular in the U.S.A. Alloy 625 is a well-established material for applications in the fields mentioned above where the corrosive conditions are somewhat less critical and more in the oxidizing range. Alloy 625 is a well-proven material for flue gas desulfurization applications (Heubner et al., 1987) and for waste incinerator scrubbers (Krause, 1988) in places which are more corrosive than the austenitic 6% Mo steels can withstand, and which do not require still higher alloyed materials, e.g., alloy C-276 or alloy 59. Alloy 625 serves well in marine applications, is well suited for overlay welding (Heubner et al., 1990), and has become the most common filler metal for welding austenitic 6% Mo steels due to the high resistance to localized corrosion of the ascast microstructure of its weldments (Heubner et al., 1989b). In many applications the high cost of the corrosion-resistant nickel-chromium-molybdenum alloys can be substantially reduced by mak-

368

7 Nickel-Based Alloys

ing use of nickel-chromium-molybdenum alloy clad carbon steel (Kirchheiner and Stenner, 1993). 7.2.2 Nickel-Based Alloys for High-Temperature Applications 7.2.2.1 Composition, Structure, Mechanical Properties, and Resulting Applications

Main Alloying Constituents Heat-resistant nickel-based alloys are suitable for applications in corroding gas atmospheres at temperatures above 550 °C, but without very strict requirements on the strength. The most important commercial, heat-resistant nickel-based alloys are shown in Table 7-7. They are essentially nickel-chromium-iron alloys with some additions of aluminum and titanium and, in the case of alloy DS and alloy 45 TM, silicon. In addition, alloy 45 TM is alloyed with small amounts of rare earth elements (RE), usually added as so-called mischmetal with cerium as its main constituent. Their carbon content is essential for high-temperature strength. They have a face-centered cubic, i.e., austenitic microstructure with precipitated carbides, the quantity of which depends on the carbon content, the final annealing temperature, and the time and temperature while in service. If high-temperature strength is required in addition to high-temperature corrosion resistance, the alloys shown in Tables 7-8 and 7-9 have to be considered. The definition of the alloys as heat-resistant or hightemperature, high-strength materials is somewhat arbitrary, but taking into consideration what is defined as a heat-resistant or a high-strength, high-temperature steel according to German standards, it

seems reasonable to define as high-temperature, high-strength materials those having a creep-rupture strength RJ105 h of above approximately 90 N/mm 2 at 600 °C, 30 N/mm 2 at 700 °C, 10 N/mm 2 at 800 °C, and 3.5 N/mm 2 at 900 °C, at least in one partial temperature range (Heubner 1995). The corresponding i?m/104 h minimum creep-rupture strength is approximately 140 N/mm 2 at 600 °C, 70 N/mm 2 at 700 °C, 35 N/mm 2 at 800 °C, 18 N/mm 2 at 900 °C, and 9 N/mm 2 at 1000 °C. The hightemperature, high-strength alloys are divided into two groups: general purpose alloys according to Table 7-8 which are basically nickel-chromium-iron alloys, and superalloys according to Table 7-9 which are nickel-chromium-based with a great variety of additions, primarily molybdenum, cobalt, aluminum, titanium, and niobium. While the main applications of the general purpose alloys are in the chemical process industry and in heat-treating and other types of furnace, the main application of superalloys is in turbine components, where the materials are at temperatures between 550 and 1050 °C. Since the iron-, cobalt-, and nickelbased superalloys are treated extensively in various monographs (Sims and Hagel, 1972; Betteridge and Heslop, 1974; Sims et al., 1987) and conference proceedings (Gell et al., 1984; Duhl et al., 1988; Loria, 1989, 1991, 1994; Antolovich et al., 1992), and in Vol. 7, Chap. 14 of this Series, this chapter only deals with the general aspects of the most common wrought nickel-based superalloys which, among the scores of superalloys, make up the lion's share and find uses in applications other than just gas turbines. Regarding the chemical composition of the high-temperature, high-strength, general purpose nickel-based alloys and of the most common nickel-based wrought su-

369

7.2 Composition, Structure, Properties, and Behavior Table 7-7. Nominal chemical compositions of heat-resistant nickel-based alloys.

Designation

Main alloying constituents (mass%)

Alloy

Trade name

Ni

Cr

C

Al

Ti

Fe

800 DS 45 TM

Nicrofer 3220 3718 So 45 TM

31 36 47

20 18 27

0.05 0.06 0.09

0.25 0.15 0.15

0.4 0.15

47 43 23

601 690 600 75

6023 6030 7216 7520

60 61 73 75

23 29 16 20

0.04 0.01 0.07 0.11

1.4

0.4

0.2 0.2

0.2 0.4

14 9 9 3

Other

2.2 Si 2.7 Si 0.10 RE

Table 7-8. Nominal chemical compositions of high-strength, high-temperature, general purpose nickel-based alloys. Designation

Main alloying constituents (mass%)

Alloy

Trade name

Ni

Cr

C

Al

Ti

Fe

800 H 800 HP AC 66 601 H 602 CA 600 H

Nicrofer 3220 H 3220 HT 3228 NbCe 6023 H 6025 HT 7216 H

31 30 32 60 62 74

21 21 28 23 25 16

0.07 0.09 0.05 0.06 0.18 0.07

0.25 0.5 0.01 1.4 2.0 0.2

0.35 0.5

47 47 39 14 9.5 9

0.5 0.15 0.2

Other

0.75 Nb

0.07 Ce

0.1 Y

0.1 Zr

Table 7-9. Nominal chemical compositions of high-strength, high-temperature nickel-based wrought superalloys. Main alloying constituents (mass%)

Designation Alloy

Trade name

Ni

Cr

Mo

C

Al

Ti

Fe

Other

Nicrofer Solid-solution and carbide precipitation-hardened:

333

4626 MoW

46

25

3

0.05

0.1

0.2

17

X 617 625 H

4722 Co 5520 Co 6022 hMo

48 54 63

22 22 22

9 9 9

0.07 0.06 0.04

1 0.1

0.5 0.2

18 1 2

3W 3 Co 1 Si 1 Co 12 Co 3.8 Nb

Age-hardenable: C-263 5120 CoTi 718 5219 Nb X-750 7016 TiNb 7520 Ti 80a

51 53 72 75

20 19 16 20

6 3

0.06 0.05 0.05 0.06

0.5 0.6 0.6 1.4

2.1 0.9 2.7 2.3

0.5 18 7.5

20 Co 5.2 Nb 1 Nb

370

7 Nickel-Based Alloys

peralloys given in Tables 7-8 and 7-9, the basic phase diagrams to be considered are either nickel-chromium-iron (Rivlin and Raynor, 1980) or, e.g., nickel-cobaltchromium-molybdenum. Indeed, in the case of a more significant departure from the nickel-chromium-iron ternary system, perhaps due to larger additions of elements such as cobalt, molybdenum, and tungsten for the purpose of increasing the high-temperature strength, the formation of sigma (a) phase and other hard intermetallic phases such as mu (JI), chi, pi, and Laves must be prevented. Such phases reduce the ductility and take up the elements added for the purpose of increasing the strength or corrosion resistance of the alloy matrix. When present in the quaternary systems (Fig. 7-13 shows an example), these phases separate the face-centered cubic austenite from the body-centered cubic chromium-molybdenum phase p. As mentioned above for the nickel-chromiummolybdenum phase diagram, it is possible to evaluate the influence of additional alloying elements on the extent of the yphase by phase computation estimates (Morinaga et al., 1984).

Co

Mo

Figure 7-13. Nickel-cobalt-chromium-molybdenum quaternary system at about 1200 °C (Woodyatt et al., 1966).

Based on the aforementioned phase diagrams, the structure of the high-temperature, high-strength, general purpose nickel-based alloys (Table 7-8) and of the nickel-based solid-solution and carbide precipitation-hardened wrought superalloys (Table 7-9) is the same, i.e., face-centered cubic with a dispersion of carbides. While the heat-resistant nickel-based alloys receive either a final soft-annealing treatment or a solution treatment, the hightemperature, high-strength nickel-based alloys and the nickel-based solid-solution and carbide precipitation-hardened superalloys get final solution-annealing treatments according to Table 7-10. The main purpose of these final solution-annealing treatments is to establish an average grain size of between approximately 100 and 200 Jim, which provide the best conditions for creep resistance in these materials (Brill, 1995). There are exceptions, e.g., alloy 602 CA remains fine-grained with a grain diameter of 40-60 |xm during final solution-annealing (Brill, 1992; Brill and Agarwal, 1993), thus providing high creep strength and improved low cycle fatigue at the same time. In this case as with other superalloys (see Fig. 7-16 below), in order to restrict grain growth only a partial solution-anneal is applied. At the same time, the carbides present dissolve to a greater or lesser extent and will precipitate again as soon as the material goes into high-temperature service. Indeed, increasing the creep resistance by means of carbide hardening is of the greatest significance for these austenitic alloys. This is because solid-solution hardening using elements like molybdenum and tungsten is only possible to a limited extent, since the precipitation of brittle intermetallic phases has to be avoided (see Fig. 7-13). An increase in the nickel content causes a decrease in the carbon solu-

7.2 Composition, Structure, Properties, and Behavior

Table 7-10. Heat-treatment guidelines for soft- and solution-annealing of nickel-based alloys for hightemperature applications a> b. Designation Alloy

Trade name

800 800 H 800 HP AC 66 DS 45 TM 601 601 H 602 CA 600 600 H 75 333 X 617 625 H C-263 718

Nicrofer 3220 3220 H 3220 HT 3228 NbCe 3718 45 TM 6023 6023 H 6025 HT 7216 7216 H 7520 4626 MoW 4722 Co 5520 Co 6022 hMo 5120 CoTi 5219 Nb

X-750

7016 TiNb

80a

7520 Ti

Softannealing

Solutionannealing

Temperature (°C)

900-980 -

1100-1150 1150-1200 1150-1200 1120-1180 1050-1100 1160-1180 920-980 1100-1190 1180-1220 _ 920-1000 1080-1150 1050 1170-1180 1160-1180 _ 1100-1180 1100-1180 1140-1160d 1) 930-1010; 2) 1040-1065 c'd 1)1150; 2) 850 c'd 1) & 2) 1080; 3)115Oc'd

a

For time at temperature see specifications, suppliers recommendations or Heubner (1987 b). Cooling is in most cases in water, whereas for thin sections air cooling might be sufficient. In bright-annealing of strip the protective atmosphere provides a sufficient amount of cooling in most cases; b Heubner (1987 b), data from Krupp VDM GmbH; c to be selected according to application and specification; d solutionannealing followed by subsequent age-hardening treatment according to Table 7-11.

bility which, for a given carbon content, leads to increased carbide precipitation, mainly at the grain boundaries. Small, equiaxed, nonagglomerating carbides are most suitable for stabilizing the grain boundaries, i.e., inhibiting grain-boundary sliding during creep. They are provided

371

principally by primary precipitated MC and M 6 C carbides, based on molybdenum, tungsten, titanium, and niobium. In view of the fact that the alloys under consideration have high chromium contents (see Tables 7-7, 7-8, and 7-9) for corrosion resistance reasons, the formation of M 2 3 C 6 carbides is inevitable. Furthermore, the MC and M 6 C carbides tend to transform into M 2 3 C 6 under long term high-temperature loading. This type of carbide has a tendency to form agglomerated grainboundary precipitates. Since carbides are brittle, these agglomerated grain-boundary precipitates provide a convenient path for crack propagation, and can thus lead to a drop in the creep resistance. On the other hand, where the precipitated carbides are equiaxed rather than agglomerated, the stresses can be reduced by grain-boundary slippage and the formation of microcracks can be prevented or at least retarded. Accompanying Elements The accompanying elements under consideration (Holt and Wallace, 1976; McLean and Strang, 1984) occur either as unavoidable components or as deliberate alloying additions, and are generally soluble in the matrix to only an extremely limited extent. They therefore tend to segregate mainly at the grain boundaries and impair their stability under creep loading. This can have both positive and negative consequences for the creep resistance. Little is known regarding the actions of the trace elements Pb, Bi, Sb, Sn, As, Te, Ga, S, P, and Ag, which have a negative effect on the creep resistance. It is believed that low melting point films form on the grain boundaries at high temperatures and cause embrittlement. On the other hand, decohesion processes are assumed to be caused by excessive

372

7 Nickel-Based Alloys

additions of these elements to disrupted lattice areas in the grain boundaries. However, these elements' clearly detrimental influence on the creep-rupture strength, which is determined by creep processes, has at least been established empirically; Fig. 7-14 provides an example. The precise maximum permissible contents of these elements are, however, not known and may differ from alloy to alloy. For instance, in the case of the widely used alloy 800H, upper limits for the elements Pb and Bi of 7 and 1 g/t, respectively, have proven favorable for the creep-rupture strength. However, the question of economy is also involved when considering whether it is worthwhile to accept the substantially higher material costs arising from a further reduction in these unwanted elements, since such a reduction leads to only a slight increase in the creep-rupture strength. It should also not be forgotten that, in practice, slight improvements in the creep-rupture strength are easily overshadowed by statistical scatter. In addition to the elements recognized as detrimental, there are a number of helpful trace elements. These are the elements

10

20 30 Pb/Bi (g/t)

40

Figure 7-14. Influence of bismuth and lead on the service life of alloy 718 at 649 °C and 690 N/mm2, according to Kent (Holt and Wallace, 1976).

Zr, B, and Mg. These elements segregate at the grain boundaries in a similar manner to the impurities discussed earlier, mainly due to the difference between their atomic size and that of the matrix elements. At the boundaries they either fix detrimental impurities, such as sulfur, or are integrated into the disrupted lattice. The latter phenomenon produces consolidation of the grain boundaries, since the diffusion processes responsible for grain-boundary slippage are inhibited or retarded. Furthermore, boron and zirconium inhibit the agglomeration of M 2 3 C 6 carbides at the grain boundaries, and thus limit the danger of microcrack formation, which has already been discussed. Moreover, boron reduces the solubility of carbon at the grain boundaries. This means that the amount of carbide at the grain boundaries increases, but the formation of coarse carbides and of continuous carbide edges is impeded. Furthermore, zirconium inhibits grain growth at elevated temperatures, since the grain-boundary energy is decreased by its integration into lattice sites in the grain boundaries, so producing a decrease in the tendency towards secondary recrystallization. Figure 7-15 shows how the creep-rupture strength of alloy 601H is favorably increased by small additions of zirconium. The lifetime of furnace components can thus be extended by a factor of 2-3. However, the use of these elements should be viewed with circumspection, since a surplus of them can be extremely detrimental. Boron contents of up to 50 g/t have thus proved to be extremely favorable, while even a slight amount in excess of this limit can produce the reverse effect. There is then, for example, a danger of low melting point borides forming, leading to corresponding embrittlement of the grain boundaries (Brill, 1995).

373

7.2 Composition, Structure, Properties, and Behavior

103

Table 7-11. Some guidelines for age-hardening treatments subsequent to solution-annealing, applicable to nickel-based alloys for high-temperature applications2. Designation

Alloy 601 H (Zr)

Alloy

Age-hardening treatment

Trade name

Temperature (°C)

Nicrofer C-263 5120 CoTi 800

1000 1200 Temperature (°C)

1400

718

800

8

1) I 720 8 II 620 8 2)1 760 10 II 650 8 furnace cooling between steps I and II in both cases

5219 Nb

Figure 7-15. Comparison of the 5000 h creep-rupture strength of alloy 601 H and with and without additions of small amounts (below 0.05 mass%) of zirconium (Heubner and Hofmann, 1989). X-750 7016 TiNb

1)I 845 24 II 705 20 2) 705 20 1) or 2) according to specification

80a

1) 700 16 2)1 850 24 II 700 16 1 3)1 925 II 750 4 1), 2), or 3) according5 to application and requirements

Age-Hardening Alloy Constituents The bottom part of Table 7-9 shows the nickel-based age-hardenable superalloys which might be regarded as the most important ones with respect to the quantities used in today's technology. The age-hardening constituents are titanium in combination with aluminum for y' precipitation hardening (see Fig. 7-9) and niobium for y" precipitation age hardening. Table 7-11 gives some guidelines on how the agehardening treatments might be conducted, but the kind of previous solution-anneal (Table 7-10) is important too. This is exemplified for alloy 718 in Fig. 7-16. The phase Ni 3 Nb is precipitated directly in its orthorhombic equilibrium form in the upper temperature range, commencing at grain boundaries. In the lower temperature range, precipitation takes place via a metastable body-centered cubic intermediate y" form, which forms as finely distributed lamellae within the mixed crystal and is thus the vehicle for the precipitation sequence. The maximum possible utilization temperature of this high-strength, high-tem-

Time (h)

7520 Ti

Heubner (1987 b).

1000-

Solution annealing Ni3Nb Start

"1 CT 900 -

1

3

1

2 800-

...1. 1.700- \

ro

.

1 1

E

*~ 600 500-

r

• t

••

^

^

-

^

' ^ Precipitation

Start 1

0.1

i

1

10

100

time (h)

Figure 7-16. Isothermal time-temperature-precipitation diagram for alloy 718, applicable to hot-rolled bars in their initial condition (Muzyka, 1972).

374

7 Nickel-Based Alloys

perature nickel alloy can thus also be seen from the diagram. It can be seen that this is around 700 °C or, better still, around 650 °C, since above these temperatures the precipitation-hardening metastable y" phase transforms within finite time into the Ni 3 Nb equilibrium phase, the material thus losing its high-temperature strength. The time-temperature paths for heattreatment are also shown in Fig. 7-16. Solution-annealing is carried out, as for some other engineering materials, as a partial solution-annealing. This brings the material into the range of equiaxed and flaky Ni 3 Nb grain-boundary precipitation, which restricts grain growth and thus influences the ductility of this alloy, even under creep loading in the 650/700 °C range. Subsequent precipitation annealing is carried out somewhat above these temperatures and is concluded as a graduated treatment. In other precipitation-hardenable nickel superalloys, stabilization of the grain boundaries is effected by means of appropriate carbide precipitation. For instance, in alloy 80a the carbide Cr 7 C 3 is precipitated at the grain boundaries in an equiaxed form during solution-annealing at 1080 °C, followed by further Cr 2 3 C 6 precipitation during the course of subsequent precipitation hardening at 700 °C. Intermediate heat-treatment at medium temperatures, such as 845 °C in the case of alloy X-750, can also be applied to obtain

the requisite carbide reaction at the grain boundaries. Mechanical Properties and Resulting Applications In general, mechanical high-temperature strength is not a primary criterion when designing with heat-resistant nickelbased alloys. Therefore, as Table 7-12 shows, only a few data are available, these referring partly to alloy 75 which, due to its low creep strength, is classified as a heat-resistant alloy in this context but is considered to be a superalloy by other authors. Indeed, alloy 75 was among the first gas turbine materials developed and is still in widespread use for casings, ducts, honeycombs, and related gas turbine applications. In cases where high-temperature strength is a criterion, the heat-resistant alloys 800, 600, and 601 are solution-annealed to a larger grain size instead of only being soft-annealed (see Table 7-10), and are thus converted into the high-temperature, high-strength alloy grades 800H, 600H, and 601H, the creep-rupture strengths of which are presented in Table 7-13. In addition to alloy 800H a new grade, alloy 800HP, had been created recently, providing still higher creep strength. This is thought to be due to the somewhat greater amounts of the alloying

Table 7-12. Creep-rupture strength (Rm/105 h) of heat-resistant nickel-based alloys3. RJ105 h (N/mm2)

Designation Alloy

Trade name

690 75

Nicrofer 6030 7520

Krupp VDM GmbH (1992).

500 °C

600 °C

700 °C

800 °C

900 °C

70 63

36 22

19

9

190

375

7.2 Composition, Structure, Properties, and Behavior

Table 7-13. Creep-rupture strength (Km/105 h) of high-strength, high-temperature, general-purpose nickel-based alloys8. RJ\05 h (N/mm2)

Designation Alloy

Trade name

600 °C

700 °C

800 °C

900 °C

1000°C

800 H 800 HP AC 66 601 H 600 H

Nicrofer 3220 H 3220 HT 3228 NbCe 6023 H 7216 H

114 126 140 156 97

53 57 52 55 42

24 26 16 16.7 17.1

10.5 11.2 5 3.7 -

4.0 -

Krupp VDM GmbH (1992).

Table 7-14. Creep-rupture strength (J*m/105 h) of high-strength, high-temperature nickel-based superalloysa. K m /10 5 h (N/mm2)

Designation Alloy

Trade name

600 °C

Solid-solution and carbideprecipitation-hardened: 333 Nicrofer 120 X 185 617 190

700 °C

800° C

900 °C

1000°C

62 98 95

27 36 43

11.4 14 16

3.9 3.6 4.5

Krupp VDM GmbH (1992).

constituents carbon, aluminum, and titanium (see Table 7-8). The ASME Boiler Code, Sec. VIII, Div. 1 Code Case, 1987, indicates a higher maximum allowable design stress for alloy 800HP. The RJ105 h creep rupture strength values for this alloy in Table 7-13 have been calculated accordingly. However, the German VdTUV data sheet 434 does not allow the use of alloy 800HP grade at temperatures below 700 °C where, due to the alloys high contents of aluminum and titanium and a resulting y'precipitation, an intolerable long-term ductility loss is to be expected, as demonstrated in Fig. 7-17 (Coppolecchia et al., 1987). Table 7-14 gives the i?m/105 h creep-rupture strengths of the high-strength, hightemperature nickel-based superalloys,

0.6

0.8

1.0

1.2

1.4

ETi+Al (mass%)

Figure 7-17. Creep-rupture elongation of alloy 800 H/ HP at 650 °C, previously annealed at 1100°C or (symbols in parentheses) 1150 °C (Coppolecchia et al, 1987).

376

7 Nickel-Based Alloys

Table 7-15. Creep-rupture strength (Rm/104 h) of high-strength, high-temperature nickel-based superalloysa. # m /10 4 h (N/mm2)

Designation Alloy

600 °C

Trade name Nicrofer

700 °C

800 °C

900 °C

1000°C

17.4 31 30

6.7 7.5 10

Solid-solution and carbideprecipitation-hardened: 180 4626 MoW 333 X 4722 Co 223 260 617 5520 Co

85 122 123

38 59 65

Age-hardenable: 718 5219 Nb X-750 7016 TiNb 7520 Ti 80a

220 270 220

60 70

650 500 460

Krupp VDM GmbH (1992). Table 7-16. Comparison of the creep-rupture strength of various wrought nickel-based alloys. Alloy group

M0+V2 W (mass%)

14 h (N/mm2)

Alloy

I5 h (N/mm2)

800 °C

900 °C

800 °C

900 °C

17 12 9

24 14 15

10.5 6 5

1

0

800 H 601 H 600 H

37 26 25

2

4.5

333

40

18

30

9

9

X 617 625 H

60 65 65

26 30 28

43 45

16 18

3

whereas Table 7-15 provides the corresponding 104 h data. Whereas for general purpose applications the 105 h creep data have to be considered in most cases, shorter lifetimes will be taken into account in gas turbine and other highly demanding special applications. Extrapolation parameters (Larson-Miller) have to be determined experimentally in order to make reliable statements. Having done so, the creep behavior can be predicted within a temperature range of 750-1000 °C for alloy 617 and 730-900 °C for alloy 800H, even if the extrapolation ratio amounts to 10 (Drefahl etal., 1986). Below 750 or 730 °C respectively, /-precipitation inter-

feres with the creep process and the extrapolation parameters will thus change with time. In Table 7-16 the i?m/104 h and RJ105 h creep-rupture strength data of nickelbased general purpose alloys and of nickelbased superalloys are compared in order to evaluate the effect of higher alloying on the creep strength. No molybdenum is added to the general purpose alloys (group 1), whereas the superalloys contain nominally 9 mass% molybdenum (group 3). Alloy 333 has 3 mass% molybdenum and 3 mass% tungsten, the latter having been taken into account as only 1.5% in order to compensate for the nearly double atom-

377

7.2 Composition, Structure, Properties, and Behavior

ic weight. As is apparent from Table 7-16, the RJ10* h creep-rupture strength of alloy group 3 with 9 % molybdenum at 800 and 900 °C is, on average, more than twice the creep-rupture strength of alloy group 1 without molybdenum as an alloying constituent. In doing such a comparison, the scatter of the creep data has to be taken into account (see below). Nevertheless, within group 1 the excellent creep strength data of alloy 800H obviously make it difficult to justify the use of the other group 1 alloys, and even of alloy 333, except in cases where the alloy's special corrosion behavior is the dominant criterion. On the other hand, the data within group 3 demonstrate that the high cobalt content of alloy 617 (Table 7-9) does not contribute to a considerably higher creep-rupture strength at 800 and 900 °C when compared to alloy 625H. The broad scatterband shown in Fig. 718 for the creep-rupture strength data of alloy 625 at 900 °C (Brill et al, 1991) is at least partly due to the fact that it comprises the creep data of both an example of alloytype 625 with a low-carbon content of 0.025 mass% and an example of alloy-type 625H with a carbon content of 0.045 mass%. In addition there is a large difference in the grain size. Within the scatterband, alloy 625H has the highest values and alloy 625 has the lowest. Data from published literature are inbetween. This shows the extent to which the influence of grain size and carbon content on the creep strength increases with time and temperature (Heubner and Kohler, 1994). The width of the creep-rupture strength scatterband also increases with time and temperature for the other high-temperature alloys, e.g., alloy X, the scatterbands for which are shown in Fig. 7-19 (Brill et al., 1991a). As shown above, this general increase in the scatterband width with time

^ 625 H; 0.045%C,3.8%Nb,1160°C final solutionanneal,100p,m mean grain diameter

± 10%

±19%

±56% 625: 0.025%C ,3.4%Nb,1120°C final solution-anneal,50|am mean grain diameter^

2

5 1(T Time (h)

2

105

Figure 7-18. Scatter-band evaluation of the creeprupture strength data of alloy 625 including data for alloy 625 H. The scatter-band comprises mean values of various melts of different origins (Brill et al., 1991 a).

±16%

10

Figure 7-19. Scatter-band evaluation of the creeprupture strength data of alloy X. The scatter band comprises mean values of various melts of different origins (Brill et al., 1991a).

and temperature is related to the increasing effect of differences in the microstructure, i.e., grain size, quantity, morphology, and distribution of carbides, and also to the increasing influence of corrosion,

378

7 Nickel-Based Alloys

mainly oxidation in air; this is presumed to differ from alloy to alloy. Future research on high-temperature creep has to clearly separate the influence corrosion might have on high-temperature creep. Despite its high creep strength, the use of alloy 625H in high temperature applications is restricted by the alloy's considerable loss of ductility over long service times at between 500 and 1000°C. The low-carbon version alloy 625 behaves somewhat better (Kohler, 1991) and may therefore find some high-temperature use, e.g., in bellow applications. However, when pressure vessel considerations come into play, other materials have to be used, e.g., alloy 617, which is, according to German VdTUV sheet 485, allowed for pressure vessel applications in the temperature range up to 1050 °C. From the standpoint of creep strength, the high-temperature, high-strength solidsolution and carbide precipitation-hardened nickel-based alloys have their application potential in the temperature range above about 650 °C. Below this either the high-temperature, high-strength stainless steels will be applied, due to their lower cost, or, in cases with higher strength demands, the precipitation-hardenable superalloys. 7.2.2.2 Corrosion Behavior, Composition, and Resulting Special Uses

Oxidation According to Table 7-7, the most important commercial, heat-resistant nickelbased alloys are essentially made up of nickel, chromium, and iron with chromium contents of 16-30%. Indeed, among the three elements aluminum, silicon, and chromium, which are able to build up pro-

tective oxide layers, chromium is the most suitable since its solubility in both nickel and iron is large and its tendency to form intermetallic phases is the smallest. Therefore in Table 7-7 only the alloys DS and 45 TM show silicon as an alloying constituent and only alloys 601H and 602 CA (Table 7-8) have appreciable amounts of aluminum. Silicon is thought to form thin oxide films very rapidly, especially in the initial stages of oxidation (Brill, 1995). Aluminum is also added in order to obtain a protective oxide film at temperatures higher than about 1000 °C, where chromium oxide evaporates noticeably, particularly in flowing atmospheres. The aluminum content of alloy 602 C A is just sufficient for the formation of a continuous film of aluminum oxide (Brill, 1992); it has to be limited in order to avoid loss of ductility due to precipitation hardening. The adhesive power of protective oxide surface layers is improved by additions of rare earth elements or yttrium to the alloy. This is why rare earth elements are added to alloy 45 TM and yttrium to alloy 602 CA (see Tables 7-7 and 7-8). Nickel-based, high-temperature alloys have been shown to generally have better resistance to oxidation under cyclic temperature loading than iron-based austenitic alloys (Brill, 1995). This general rule is not always well observed, since besides iron and nickel the remainder of the alloying constituents have to be taken into account. Furthermore, the oxidation resistance as a materials behavior, and not a materials property, will also depend on the surrounding medium, i.e., the test or service conditions. Here cyclic or interrupted test conditions have been shown to be much more critical than continuous service. Therefore the evaluation of materials is preferentially done under alternating conditions, e.g., cycling between 16 h exposure

7.2 Composition, Structure, Properties, and Behavior

to air at the test temperature followed by subsequent cooling to ambient temperatures for 8 h. Figure 7-20 shows the weight change of some nickel-based alloys under such test conditions when exposed to air and cycled between 1000 °C and ambient temperatures. After a total of about 1000 h, the alloys X, 601H, and 617 exhibited the smallest weight change. If the intergranular penetration of the oxidation (as indicated in Table 7-17) is considered as well, alloys 601H and 617 come out best, whereas alloy 800H would not be useful under these conditions. This corresponds well with practical experience, according to which alloy 601H is selected as a construction material if excellent oxidation resistance is required up to about 11 SOX, e.g., for the design of advanced gas-fired, heat-treating furnaces where radiant tubes, made from alloy 601H, run at about 10501150 °C (about 1290-2100 °F). Under considerations of cost alloy 617, due to its alloying constituents of cobalt and molybdenum (Table 7-9), will only be selected for service under oxidizing conditions at Table 7-17. Corrosion of nickel-based alloys after alternating exposure to air at 1000°C, the time of exposure being made up of cycles where every cycle comprises 16 h at 1000 °C and 8 h at ambient temperatures a Alloy

800 H 333 X 617 625 H 601 H a

Material Time of trade exposure name (h)

Nicrofer 3220 H 4626 MoW 4722 Co 5520 Co 6022 hMo 6023 H

Brill etal. (1991b).

1056 744 1056 1056 408 1056

Weight Additional change intercrystal(g/m2h) line deterioration (mm)

-0.03 -0.05 + 0.004 + 0.02 -0.1 + 0.066

0.2 0.06 0.12 0.06 0.07 0.06

379

JI-301200 900 600 Time (h) Figure 7-20. Weight change of nickel-based hightemperature alloys during continuous repeated cyclic exposure to air for 16 h at 1000 °C followed by cooling for 8 h to ambient temperatures (Brill et al, 1991 b).

10 800

1000

1200

1400

Temperature (°C)

Figure 7-21. #m/5000 h creep-rupture strength of alloys 601 H and 617 at very high temperatures (Drefahl and Hofmann, 1986).

about 1000 °C if its higher creep strength is the prevailing design criterion. For service above 1000 °C, the creep-rupture strengths of the alloys become closer together the higher the temperature. This is demonstrated in Fig. 7-21. Alloy 602 CA is a recent alloy development and is superior to alloy 601H with respect to both creep strength and oxidation resistance up to 1200 °C (Agarwal and Brill, 1993 a; Brill 1992). Alloy X, despite ranking third according to the results presented in Fig. 7-20 and Table 7-17 has a firm position as a well-proven material for the construction of combustion chambers

380

7 Nickel-Based Alloys

in gas turbines. It obviously exhibits a well-balanced combination of creep strength, oxidation resistance, and materials cost, which makes it a favorite material for this application, whereas alloys 800H and 333 are predominantly used at somewhat lower temperatures.

300

200-

100-

20

30

Mass% Fe

Carburization Carburization is observed in several applications of high-temperature alloys in oxidizing and carburizing environments. Carbon is transferred by gaseous molecules such as CO, CO 2 , CH 4 , and other hydrocarbons into the metal, and internal carbides M 2 3 C 6 and M 7 C 3 (M = Cr, Fe, Ni) precipitate as a result. This may cause a loss in ductility of the material. In most cases the ingress of carbon is retarded by the presence of an oxide layer. Therefore high chromium contents may promote the resistance to carburization, where diffusion-inhibiting chromium oxide layers can be generated on the surface of the metal. However, these are only effective as long as they remain sealed. Cracked oxide layers may thus lead to increased carburization. On the other hand, increasing nickel contents are known to reduce the alloy's susceptibility to carburization (Kofstad, 1988). Figure 7-22 shows the weight gain of some nickel-based alloys after cyclic exposure to a carburizing environment (CH 4 /H 2 atmosphere with activity of carbon ac = 0.8) at 1000 °C. The total time of 864 h was made up of 24 h cycles, each cycle consisting of 16 h at 1000 °C and 8 h at ambient temperatures. The diagram shows the weight gain by carbon pick-up to clearly be a function of the alloys' iron content. But there are two exceptions. The carburization of alloy 30 is lower than would be expected from its iron content. Alloy 30 as

Figure 7-22. Weight gain by carburization of some nickel-based alloys after cyclic exposure to a carburizing atmosphere (CH 4 /H 2 , ac = 0.8) at 1000°C. The total time of 864 h is made up of 24 h cycles, each cycle consisting of 16 h at 1000 °C and 8 h at ambient temperature (Brill et al., 1991 b).

the heating resistance alloy version of alloy 800 contains nominally 2.5 mass% silicon, and this alloying element is known to decrease the susceptibility to carburization (Kofstad, 1988). This effect is also observed for alloy 333, where the silicon content of nominally 1 % makes the weight gain by carburization comparatively low. On the other hand, the weight gain by carburization of alloy 625H is much higher than would be expected from its low iron content. Comparing alloy 625H to the other alloys, this behavior can only be attributed to the high niobium content of alloy 625H, in contrast to previous observations on the effect of niobium in other alloys. The carburization resistance of alloy 45 TM is in the same range as that of alloys 625 and 601 (Brill and Klower, 1993). In practice, the high cost of alloy 617 normally prevents its use as a carburization-resistant material, but alloys 333, 45 TM, 601, and 600 are applied. It has been shown (Heubner etal., 1991a) that the creep strength of alloy 333 is not impaired by carburization, and the creep ductility was also observed to be high under carburizing conditions. With alloy 800H carburization increases the creep-rupture

7.2 Composition, Structure, Properties, and Behavior

strength considerably, whereas the creep ductility is lowered to a level which is obtained in air after long exposure times (Heubner et al., 1991a). Finally it has to be mentioned that new high-temperature Ni3Al intermetallic phase materials, which are still under development, show excellent resistance to carburization [much better than that observed for alloy 617 (Brill, 1990)]. Sulfidation In addition to oxidation and carburization, attack by sulfur-bearing media must' also be regarded as a high-temperature corrosion phenomenon. This type of corrosion is in general more deleterious than those discussed up to now, since metal sulfides have lower melting points than the corresponding oxides and carbides. This may then lead to catastrophic corrosion via liquid phases. The phenomenon is severest in materials with a high nickel content, since a low melting point eutectic alloy system (Ni-Ni 3 S 2 ) exists at temperatures as low as 650 °C. Whether sulfidation occurs, and to what extent, depends largely on the make-up of the aggressive medium. As well as thermodynamic considerations, kinetic aspects are likely to play a decisive role. In oxygen-bearing media, alloys with high chromium contents can form chromium oxide surface layers which protect the alloy from sulfidation, at least temporarily. Therefore, in order to provide for maximum resistance to sulfidation, the alloy's nickel content should be low and the chromium content high. Consequently, corrosion resistance in sulfur-bearing atmospheres is to be expected by preference from alloys like AC 66 and 45 TM, but alloy 800H has proved to be a suitable material in many atmospheres which are

381

alternately oxidizing/carburizing and contain sulfur at the same time, as may occur in coal gasification and in the petrochemical industry (Agarwal and Brill, 1993 b; Brill and Klower, 1993). Nitridation Most high-temperature materials are resistant to attack by molecular nitrogen up to approximately 1000 °C. Above this temperature, however, dissociation into atomic nitrogen commences, and this reacts with the nitride-forming alloy components chromium, iron, titanium, and aluminum. In ammonia atmospheres this reaction occurs at relatively low temperatures. The corrosion products are hard, brittle, nitride phases. The ductility of the material in the surface zone is then decreased, leading to possible cracking. Furthermore, due to the formation of nitrides, these alloying elements are no longer available for their basic function, for instance, chromium and aluminum for the formation of protective oxide layers, or titanium and aluminum for precipitation hardening. From an alloy point of view, nitrogen pick-up can, particularly under ammonia, be countered by high nickel contents and simultaneous reductions in the amounts of the elements that have an affinity to nitrogen. Up to approximately 950 °C, nickel has practically no reaction with atomic nitrogen. Accordingly, alloy 600H is to be recommended for all applications where nitridation might be a problem. The excellent behavior of this alloy in various nitriding atmospheres has been demonstrated in a recent study (Rosa and Smith, 1987). Chlorination High-nickel alloys are necessary as construction materials for halogen-containing atmospheres. The relatively good resis-

382

7 Nickel-Based Alloys

tance to corrosion possessed by these alloys can be attributed to the formation of a protective layer of nickel halides. At elevated temperatures of up to approximately 700 °C, these have considerably lower vapor pressures than, for example, chromium or iron halides. In the presence of oxygen, alloy 600H proved to be a usable material at 750 °C, as demonstrated in Fig. 723, whereas in chlorine-containing flue gases from solvent incineration at 650 °C the alloys 625 and C-276 showed superior performance (Brill etal., 1990). In exposure tests in a waste incineration atmosphere (2.5 g m~ 3 HC1, 1.3 g m " 3 SO 2 , 9% O 2 , balance N 2 ) in the temperature range 550-850 °C, alloy 45 TM gave the best performance (Agarwal and Brill, 1993 b; Brill and Klower, 1993). Hot Gas Corrosion Hot gas corrosion is only one of various types of deposit-induced corrosion which may be observed with high-temperature nickel-based alloys (Heubner and Hofmann, 1989). It is due to the condensation of sodium sulfate on sites which are below a temperature of about 1000 °C. The resistance to sulfate corrosion can be improved

600 H

ERNiCr-3 -600 H "C-4 600 H+600 =

-TOO Ni99.6 -200 200

400 Time (h)

600

800

Figure 7-23. Weight change by chlorination of some nickel-based alloys after cyclic exposure to a chlorinating atmosphere (Ar/0.25% C12/2O% O2) at 750 °C. The time indicated is made up of 24 h cycles, each cycle consisting of 16 h at 750 °C and 8 h at ambient temperatures (Brill et al, 1991 b).

by increasing the chromium content of the metallic material, but the chromium content of current wrought metallic materials is limited to about 30%. It is therefore in such a case recommended that the design or the service conditions of the combustion system be changed in order to avoid the contact of alkaline sulfate with metallic parts at temperatures below 1000 °C. 7.2.3 Nickel-Based Alloys for Electrotechnical Applications 7.2.3.1 Heating Resistance Alloys Heating resistance alloys are designed to transform electrical energy into heat. Therefore heating resistance alloys constitute a special species of heat-resistant alloys which are usually manufactured into wire or strip, and are used where electrical resistivity is required in combination with heat resistance. In this context heat resistance is, in most cases, resistance to cyclic oxidation in air, although resistance to other atmospheres might be required as well. Since heating resistance wires are supported or insulated by ceramics in many applications, they have to exhibit sufficient compatibility with these materials. Some applications also require a certain amount of creep resistance. Other requirements are sufficient cold formability and weldability. Table 7-18 shows the nominal chemical compositions of commercial nickel-based heating resistance alloys. Alloys 80 and 70 are nickel-chromium alloys, whereas alloys 60, 45, and 30 are based on the nickel-chromium-iron ternary system. On comparison with the heat-resistant nickel-based alloys in Table 7-7, it becomes obvious that two of the heating resistance alloys are closely related to the heat-resistant alloys, i.e., alloy 80 to alloy 75 and alloy 30 to alloy 800. Howev-

7.2 Composition, Structure, Properties, and Behavior

383

Table 7-18. Nominal chemical composition of commercial nickel-based alloys for heating resistance applications. ]Designation

Alloy

Trade name

Main alloying constituents (mass%) Ni

Cr

Fe

Si

C

RE a

Other b

80 70 60

Cronix 80 Cronix 70 Cronifer II

78 68 60

20 30 15

0.4 0.5 22

1.2 1.2 1.5

0.04 0.04 0.05

0.04 0.04 b -

45 30 30 Special

Cronifer 45 Cronifer III Cronifer III So

46 31 32

23 20 20

28 46 44

1.8 2.5 1.9

0.03 0.05 0.05

0.03 0.08 0.27

0.04 Ca c 0.06 Zr

a

RE: rare earth elements; b REs are at about 0.015% when the alloy is manufactured as strip; c 0.02% RE may be substituted for Ca and Zr.

er, the difference is that all the heating resistance alloys contain appreciable amounts of silicon. Besides its effect on retarding carburization (see Sect. 7.2.2.2), silicon is known to improve the protective effect of the chromium oxide layer when the alloys undergo intermittent service (Pfeiffer and Thomas, 1963). Other important alloy additions are rare earth elements (RE), which are most frequently added as mischmetal where usually cerium, lanthanum, and e.g., neodymium are the main components. One of the alloys mentioned in Table 7-18 contains additions of calcium and zirconium instead of rare earth metals. All these oxygenactive elements improve the oxidation resistance and particularly the adherence of the protective chromium oxide scale during intermittent high-temperature service (Pfeiffer and Thomas, 1963; Kofstad, 1988). So the nickel-based heating resistance alloys are designed to have a facecentered cubic microstructure with reactive elements and carbides dispersed in the metallic matrix. The carbon content will influence the creep strength in the same way as it does in high-temperature, highstrength nickel-based alloys, but the alloys

will be annealed in most cases to smaller grain sizes in order to achieve good coldforming characteristics. The creep strength in terms of stress for 1% strain after 1000 h i? pl /10 3 h at 1000 °C of the nickel-based heating resistance alloys is about 4N/mm 2 with the exception of alloy 45 for which 6 N/mm 2 has been obtained (VDM Nickel-Technologie AG, 1990). With respect to this property, the nickel-based heating resistance alloys are by far superior to the ferritic iron-chromium-aluminum heating resistance alloys. Electrical resistivity as one of the paramount properties of nickel-based heating resistance alloys, as shown in Fig. 7-24. In the case of alloys 80 and 70 the temperature dependence is small, whereas it is greatest for alloy 30. The irregularities observed, especially in the resistivity-temperature curves of alloys 80, 70, and 60, are due to structural transformations and therefore, below about 550 °C, a lower resistivity will be observed after rapid cooling of the materials. Heating resistance alloys are usually required to undergo intermittent service. Therefore the lifetime will be finished

384

7 Nickel-Based Alloys

1.30 -

1E

1.25 -

AlloyVO

•-?'

E

Alloy 60

« 1.20 -

resi

-2 1.15 Alloy80 / /Alloy30

•4—

X

1 1.05 1.00 -0

200

400

600

800

1000

1200

Temperature (°C)

Figure 7-24. Electrical resistivity of commercial nickel-based heating resistance alloys as a function of temperature for the slowly cooled condition (VDM Nickel-Technologie AG, 1990).

when cyclic exposure to high-temperature corrosion, i.e., mostly oxidation, has led to the resistivity increasing to beyond a certain limit (e.g., 10%), or eventually to complete destruction. Consequently, a standard test commonly used for the evaluation of heating resistance alloys is a cyclic switch-on, switch-off test where a wire of 0.4 mm diameter, which is freely exposed to air, is rapidly heated to the test temperature, e.g., 1200 °C, by direct current, held at this temperature by means of this current for 2min, and subsequently cooled down to ambient temperatures after the current has simply been switched off. The test temperature is controlled by means of optical photometry. This cycle is repeated until the electrical resistivity at the test temperature has increased by a certain amount e.g., 10%, or until the wire is burnt through, depending on what has been determined as the lifetime limiting factor in the practical application under consideration.

Typical lifetimes until through-burning in such a test when conducted at a temperature of 1150 °C may be about 4900, 2900, and 1600 cycles for alloys 80, 60, and 30, respectively. These lifetime values exemplify the different potentials of today's commercial heating resistance nickel-based alloys, but not every application requires the highest lifetime according to this type of test. As is to be expected, the average life cycle values decrease with test temperature, and the maximum permissible service temperatures are different for each type of alloy, as shown in Table 7-19. Selection of the appropriate alloy type depends on the service conditions, i.e., the service temperature, furnace atmosphere, and design of the heating equipment (VDM Nickel-Technologie AG, 1990). In oxidizing atmospheres the nickel-based alloys are unsuitable above 1150°C because the oxidation rate is too high, and therefore iron-chromium-aluminum alloys are recommended. If resistance to sulfidation is required, the use of alloys which are low in nickel or of iron-chromium-aluminum heating resistance alloys is recommended. If resistance to chlorination is required, e.g., in glazing furnaces for ceramic hardware or for resistance to water vapor-containing atmospheres, the high-nickel alloys

Table 7-19. Maximum permissible service temperature of nickel-based heating resistance alloy wire (>2mm diameter)3. Designation Alloy

Trade name

80 70 60 45 30

Cronix 80 Cronix 70 Cronifer II Cronifer 45 Cronifer III

Max. perm, service temperature (°C) 1230 1250 1150 1170 1100

VDM Nickel-Technologie AG (1990).

385

7.2 Composition, Structure, Properties, and Behavior

are superior. Nickel-based heating resistance alloys are also preferred to ironchromium-, aluminum alloys in the case of nitrogen atmospheres. Certain applications, e.g., in heat-treating furnaces, may require resistance to alternating oxidizing and reducing atmospheres which rapidly decompose alloys 80 and 60 by internal oxidation of the chromium preferentially along the grain boundaries: so-called green rot. Alloy 70 has better potential to resist this type of damage and is generally to be recommended for use in low-oxygen atmospheres. Finally, within the scope of this section, the copper 44mass% nickel-1 mass % manganese alloy type has to be mentioned, which is known for its low coefficient of electrical resistivity between ambient temperatures and about 100°C, and has the trade name Konstantan. Due to its high resistance to general corrosion, this alloy is also used for low-temperature heating resistance applications up to about 600 °C, e.g., in heating covers and pillows. 7.2.3.2 Spark Plug Electrode Alloys Spark plug electrodes have to withstand corrosion and spark erosion in high-temperature automotive environments. Table 7-20 gives a selection of commercial spark plug alloys which are in use today. The structure of these alloys is face-centered

cubic. Alloy 836 is high in chromium, whereas the other two alloys are not. Indeed, chromium has to be reduced if a requirement for high thermal conductivity dominates. Another aspect is the very low carbon content of the two high-nickel alloys, due to the requirement for excellent cold formability and long die life in coldstamping operations, whereas the service characteristics of the spark plug alloys have to be evaluated in direct motor tests. 7.2.33 Soft Magnetic Alloys General Survey Within the broad family of soft magnetic materials (see Vol. 3 of this Series), nickel-based alloys constitute an important group. Table 7-21 shows the nominal chemical compositions of commercial soft magnetic nickel-based alloys. Obviously iron is the most important alloying element, but in some alloys molybdenum, copper, and chromium are also added. Manganese and silicon are accompanying elements. All these alloys have a face-centered cubic microstructure, and they are ferromagnetic at ambient temperatures. As is apparent from Table 7-21, two classes of alloys have been developed: high-nickel alloys with about 76-80 mass% Ni, and a class of alloys with about 50 mass% Ni. Apart from these two classes, an alloy type

Table 7-20. Nominal chemical composition of some commercial nickel-based alloys for spark plug electrodes. Main alloying constituents (mass%)

Designation Alloy

Trade name

Ni

Cr

Fe

Mn

Si

C

Other

836 522

Nicrofer 7615 Nickel-Chrom-2Mangan Nickel-Chrom-2Mangan-Si

77 95

15 1.8

7.5

0.3 2

0.3 0.5

0.01 0.002

0.15 Zr

95

1.6

2

1.5

0.003

-

386

7 Nickel-Based Alloys

Table 7-21. Nominal chemical composition of some important commercial soft magnetic nickel-based alloys3. Designa-

Main alloying constituents (mass%)

nation

Magnifer 7904 7754 75 53 50 36 a

Ni

Fe

Mo Cu

80 77 76 55 48 36

14 13 16 44 51 63

5 4

5 5

Cr

2

Mn Si

0.5 0.5 0.6 0.4 0.4 0.3

0.3 0.3 0.2 0.2 0.15 0.2

The carbon content is about 0.02 mass%.

-200

disordered

E

ordered

\

-2 -3 -

with 36% Ni also exists. This situation is a consequence of magnetic anisotropy energies and saturation polarization in the iron-nickel system, both governing the magnetic properties. In its central part Fig. 7-25 shows the magnetocrystalline anisotropy constant K± between 30 and 100% nickel. In the as-quenched condition Kt is positive up to about 75 % nickel and adopts negative values at higher nickel contents, depending on whether <100> or <111> is the direction of easy magnetization. As the diagram shows, in the 75 % nickel region the cooling velocity has a large influence, with slowly cooled alloys exhibiting much lower Kx values than alloys in the as-quenched condition. This is the way that K± depends on the degree of atomic ordering which is achieved during slow cooling or low-temperature annealing. In contrast to the magnetocrystalline anisotropy, the magnetostrictive anisotropy in the nickel-iron system is less influenced by the cooling rate, as demonstrated in the lower part of Fig. 7-25.

+30 + 20 ordered

0 -

-10 40

60

100

Mass% nickel

Figure 7-25. Saturation polarization J s , Curie temperature Tc, magnetocrystalline anisotropy constant Kl9 uniaxial anisotropy Ku (after 450°C tempering treatment), and magnetostrictive anisotropy A i n and A100 of binary iron-nickel alloys. K1 and X are indicated for the as-quenched (disordered) and for the slowly cooled or annealed (ordered) condition, taken from Pfeifer and Radeloff (1980).

High-Nickel Alloys with 76-80 mass% Nickel When looking for greatest ease of magnetization in terms of highest permeabilities, both the magnetocrystalline anisotropy constant Kx and the magnetostrictive anisotropy constants A 111? A 100 , and Xs (average magnetostrictive anisotropy con-

387

7.2 Composition, Structure, Properties, and Behavior

stant) have to be zero. From Fig. 7-25 it would be expected that this could best be achieved with alloys containing about 75-80 % nickel. Although it is not possible to make all the magnetostrictive constants zero at the same time, good results are obtained if X1X1 or Xs together with K± are brought near to zero. This is mainly done by alloying with molybdenum or copper with a suitable heat-treatment to provide suitable ordered conditions. Figure 7-26 shows the zero position of K± and the AX11 and As constants in the N i - F e - M o - C u system. High initial permeability is obtained at the intersection of the lines K± = 0 and As = 0. In the diagram additions of copper of 5, 10, and 14 mass% replace the same amounts of iron. When the final annealing treatment is done at 550 °C, rapid cooling has to be applied, otherwise the K± = 0 line will be shifted towards higher molybdenum contents, as shown by the position of the K1 = 0 line established for the final 480 °C tempering treatment. Comparison of Fig. 7-26 with Table 7-21 shows the 80 Ni-5Mo alloy and the 77Ni-4Mo-5Cu alloy to be close to the positions K± = 0 / A n i = 0/515 °C final tempering temperature and K1 = 0/Xs = 0/ 480 °C final tempering temperature, respectively. Chromium may replace molybdenum as in the case of the 77Ni-5Cu2Cr alloy. Figure 7-27 exemplifies the permeabilities of the 80Ni-5Mo soft magnetic alloy as obtained on an industrial 30 t melt of Magnifer 7904. The material had been melted in an electric arc furnace and, after a vacuum oxygen decarburization treatment, cast into ingots which had been hot and subsequently cold rolled to strip of 0.07 mm thickness. The strip was manufactured into toroidal tape-wound cores. After a final anneal at 1200°C in a dry hydrogen atmosphere and a slow furnace cooling at a rate of 0.9°C/min, the cores

85

90 Mass%

75

70

Ni

Figure 7-26. Zero position of the magnetocrystalline anisotropy constant Kl9 the magnetostrictive constant l n i , and the average magnetostrictive constant As in the N i - F e - M o - C u system (Pfeifer and Radeloff, 1980).

500000

100000 460

480

500

520

Take-out temperature (°C)

Figure 7-27. Permeabilities at 20 °C of one heat of the 80Ni-5Mo soft magnetic nickel-iron alloy Magnifer 7904 after being manufactured into toroidal tapewound cores, with final annealing at 1200°C in a dry hydrogen atmosphere, furnace cooling at a rate of 0.9 °C/min, taken from the furnace at the temperature indicated, and further cooled down in air (Hattendorf, 1991).

were taken from the furnace at the temperatures indicated in the diagram and allowed to further cool down in air. As the diagram shows both the initial permeability /i 4 , measured at 4 mA/cm, and the maximum permeability fimax increase when the take-out temperature goes up to 480/ 490 °C. Obviously industrial annealing furnaces for soft magnetic materials of this kind have to work very precisely in order

388

7 Nickel-Based Alloys

to obtain the desired high permeabilities for a whole furnace load. At the position of greatest initial permeability in Fig. 7-27, the magnetocrystalline anisotropy constant ^ = 0 , whereas K± < 0 at lower take-out temperatures or with slower cooling before take-out. In this case <111> is the preferred direction of magnetization. At the same time the composition of the alloy also caused the magnetostrictive anisotropy constant lil± to be close to zero. This avoids greater dimensional changes during magnetization, i.e., low internal stress which results in a rectangular hysteresis loop, as indicated in Fig. 7-27. At higher take-out temperatures or with more rapid cooling, K1>0 and <100> is the preferred dirction of magnetization. Since X100 i=- 0 dimensional changes will occur, resulting in more internal stress. This is the reason why the permeability decreases more rapidly at higher take-out temperatures; the hysteresis loop will also be flatter, as shown in Fig. 7.27. For many applications a small variation in the permeability as a function of the service temperature is required. The variation of the initial permeability /i 4 with the service temperature between —20 and + 80°C is shown for one melt of the 80Ni-5Mo alloy Magnifer 7904 in Fig. 7-28. The cores taken out at 495 °C show a small variation in the initial permeability with a service temperature of between 20 and 80 °C, but a sharp drop if the service temperature is lowered to — 20 °C. ^ = 0 at the highest point of the curve, whereas K1 > 0 at lower temperatures and K1<0 at higher service temperatures. If the takeout temperature from the furnace is lowered, the Kx = 0 position is shifted to lower service temperatures so that the flat portion of the curve is larger but with a lower permeability level at 20 °C. In practice, a compromise between high permeability at

T

-20

1

1

1

1

0 20 40 60 Service temperature (°C)

80

100

Figure 7-28. Initial permeability fi4 at various service temperatures of one melt of the 80Ni-5Mo soft magnetic nickel-iron alloy Magnifer 7904 after being manufactured into toroidal tape-wound cores, with final annealing at 1200 °C in a dry hydrogen atmosphere, furnace cooling at a rate of 0.9 °C/min, and different furnace take-out temperatures TE (Hattendorf, 1991).

20 °C and a small dependency of the permeability on the service temperature has to be made, dependent on the application requirements. Ground fault-current interruptors have to respond not only to ac (alternating current) but also to dc (direct current) pulses as a consequence of the growing use of electronic equipment in daily life. This requires a hysteresis loop with low remanence combined with high steepness and thus a high pulse permeability. This situation is illustrated in Fig. 7-29. Starting with a highly permeable alloy with Kx = 0/ X = 0 by annealing below the Curie temperature in a magnetic field (so-called magnetic annealing), with the field perpendicular to the later direction of magnetization during service, directional short-range-order with a resultant uniaxial anisotropy Ku and thus a hysteresis loop with lower remanence Br are achieved. Since the Curie temperature of the 76-80 mass% nickel alloys is at 400-420 °C, the magnetic annealing has to be done, e.g., for 1-1.5 h at 300-350 °C.

389

7.2 Composition, Structure, Properties, and Behavior Rectangular loop 0,8 T1

_B

-y0"—l

1

0,81"'

• —

f

/

50Hz

Flat loop

-

I

-J

_B

B

B

H 5 A/m

0,8T

5 A/m

0,8 T

5 A/m

-

5 A/m

b)

Figure 7-29. (a) Sinusoidal alternating magnetic flux density, and (b) pulsed magnetic flux density, and the resulting hysteresis loop without (rectangular loop) and with (flat loop) magnetic annealing of an 80Ni-5Mo nickel-iron alloy, fabricated into toroidal tape-wound cores (Hattendorf, 1991). H is the magnetic field strength, B the induction, and t the time.

After their final high-temperature annealing treatment, the soft magnetic nickel-iron alloys mentioned in Table 7-21 are also mechanically soft. The 80Ni-5Mo alloy in this condition shows a hardness of, e.g., only 90-100 HV 10. By alloying with suitable amounts of Ti and Nb and an agehardening treatment, a considerable increase in the hardness can be achieved. Since the particles causing this precipitation hardening are small in comparison to the Bloch-wall thickness, high permeabilities can be obtained at the same time. These types of alloys have proved their usefulness, e.g., as magnetic heads in tape recorders. Important applications of the soft magnetic high-nickel alloys with 76-80 mass% nickel include inductive components in integrating current transformers for current circuit breakers, instrument transformers, low frequency transducers, and relay parts and screens.

Nickel Alloys with About 50 Mass% Nickel While the low-nickel alloy type with about 36 mass% nickel, as mentioned in Table 7-21, is of interest mainly due to its high specific electrical resistance, the two alloy types with about 50 % nickel have a broad number of applications in particular because of their high saturation polarization / s (see Fig. 7-25). However the notorious crystalline anisotropy constant, Kt > 1 and thus <100> is the preferred direction of magnetization. Since Kt # 0 the permeabilities are lower than with the high-nickel alloys. However, <100> being the preferred direction of magnetization provides the opportunity of increasing the permeability by making use of textures with <100> in the rolling direction. These textures are either {100} <001> or {210} <001>. According to Fig. 7-30 a, the {100} <100> or so-called cubic texture is obtained after a high degree of cold working

390

7 Nickel-Based Alloys

1000

a)

.H"

0 900o

2 800OJ

1 700 -

Degreei -*- improving

g» 600 '"§ 1000 c ro

_a, 900 " ro

1u 800 (U

•4-

~ 700600

Normal secondary recrysrallization !

88

90

r

,

j

92 94 96 Cold rolling (%)

,—

98

Figure 7-30. Influence of cold rolling, intermediate annealing treatment, and final annealing on the microstructure of nickel-iron alloys with about 50 mass%, according to Pfeifer and Radeloff (1980).

and final annealing at temperatures between 900 and 1050 °C. Its distinctness is generally better the higher the degree of final cold working and the smaller the grain size or the lower the intermediate annealing temperature before the final cold-rolling operation. On annealing at 1100-1200 °C? a secondary recrystallization occurs, destroying the cubic texture. The range of suitable intermediate annealing and the degree of cold working at which secondary recrystallization occurs correspond to those of the cubic texture, since this is a prerequisite for secondary recrystallization. Figure 7-30 b shows the conditions under which this secondary recrystallization develops a {210} <001> texture. The border lines between the fields in Fig. 7-30 will vary to some extent from heat to heat according to the alloying elements present and to the degree of internal cleanness, i.e., the extent of in-

clusions of nonmetallic deoxidation products. While the cubic texture causes a sharp, rectangular hysteresis loop, the permeabilities are greater with the {210} <001) texture due the larger, accompanying grain size. The 50 % nickel alloys can also be provided with a uniaxial magnetic anisotropy by magnetic annealing. On annealing in a field perpendicular to the later direction of flux during service, very flat hysteresis loops can be obtained. Although the pulse permeability is lower than with the highnickel alloys, due to the higher saturation induction obtainable with the 50 % nickel alloys, the usable excursion range of induction is greater. A considerable increase in the pulse permeability at a high excursion range of induction is achieved by transverse magnetic annealing of alloys with a {210} <100> texture. While annealing in a magnetic field perpendicular to the later applied field during service will always be the first choice if the requirement of a flat loop is predominant, annealing in a magnetic field longitudinal to the later direction of service will result in higher permeabilities, since this will exclude the rotation of magnetic domains. Figure 7-31 shows what happens when an alloy with 55 % nickel (Magnifer 53 from Table 7-21) is annealed in a longitudinal magnetic field. In most cases an annealing treatment at about 440 °C is used in order to obtain both high /i 4 and /i max permeabilities with high saturation induction at the same time, as required for use with instrument transformers. On the other hand, on annealing at a temperature of, e.g., 480 °C a high pulse permeability can be obtained as a result of a low ratio of remanence BR to saturation induction Bs, which is to be aimed at for use with pulsing dc current. This behavior results from the opposite influence of the annealing temperature in the

7.2 Composition, Structure, Properties, and Behavior 100000

200000

391

Influence of Dislocations, Grain Size, and Internal Cleanness

75000 -

,50000 25000 800 -

600 400200 -

400

420

440

460

480

500

Annealing temperature (°C)

Figure 7-31. Initial permability /i 4 , maximum permeability fimax, usable excursion rate of induction B, and ratio of remanence BT to saturation induction Bs of the 55Ni-44Fe alloy Magnifer 53 manufactured into strip of 0.1 mm thickness and fabricated into toroidal tape-wound cores, high-temperature annealed at 1200°C in dry hydrogen with a dew point of — 60 °C, resulting in a {210} <001> texture, annealed in a longitudinal magnetic field until saturation, and exposed to a 50 Hz sinusoidal ac current (^4, /imax), or to a 50 Hz pulsed sinusoidal dc current (B). BJBS is from the static hysteresis loop (Hattendorf, 1991).

longitudinal magnetic field on the magnetocrystalline anisotropy constant K± and the uniaxial magnetic anisotropy Ku. Important applications of soft magnetic nickel-iron alloys with about 50 mass% nickel include inductive components in instrument transformers, LF power transducers, chokes (coils), and integrating current transformers for current circuit breakers and relay parts.

It has been known for a long time that the magnetic properties of soft magnetic nickel-iron materials are not only determined by the magnitudes of anisotropy, but also to a very large degree by any disturbance of the crystal lattice and the microstructure (Pfeifer & Radeloff, 1980). Such disturbances are dislocations, grain boundaries, and nonmagnetic inclusions. Dislocations have to be carefully minimized by making sure that no deformation occurs after the last high-temperature anneal. This excludes, e.g., any quenching operation, and calls for slow cooling and for great care during subsequent handling and assembling of parts made from these soft magnetic alloys. Grain boundaries have to be minimized by aiming at a large grain size during the final high-temperature anneal which is therefore best done for some hours at 1100-1200 °C. Nonmagnetic inclusions result from the deoxidation treatment of the melt before casting into the molds. Depending on the type of deoxidation treatment applied, these inclusions can be oxides of Al, Mg, Ca, etc. In 30 t scale meltshop practice, their average diameter has been found to be about 1 jim which corresponds well with the Blochwall thickness in the high-nickel-iron alloys. The smaller diameter fractions of the inclusions are not too far from the 0.2 jam Bloch-wall thickness of the 50% nickeliron alloys. Hence, in both alloy types, an interaction with the Bloch-walls is to be expected, so hampering their movement. However, these inclusions also handicap the grain growth during high-temperature annealing so that, in practice, it is difficult to separate both factors. Figure 7-32 shows an evaluation of the initial permeability as function of the grain

392

7 Nickel-Based Alloys

will handicap the grain-boundary movement.

400000 -

S 300000-

. 200000 -

s 100000

80

7.2.3.4 Controlled Expansion Alloys

10 140 100 120 Average grain size (^jrn)

160

Figure 7-32. Initial permeability //4 as function of average grain diameter for various heats of the 80Ni5Mo soft magnetic nickel-iron Magnifer 7904. The materials were manufactured into strips of 0.07 mm thickness, fabricated into toroidal tape-wound cores, and annealed for 6 h at 1100°C or for 4 h at 1200°C (Hattendorf, 1991).

size for a number of different melts of the 80Ni-5Mo alloy Magnifer 7904. When examined for their internal cleanness according to Japanese Industrial Standard JIS G 0555, the materials were divided into three classes: cleanness lower than 5, between 5 and less than 10, and 10 or above, the cleanness numbers being the inverse volume fraction (%) of nonmetallic inclusions in the alloy. From Fig. 7-32 it is evident that the initial permeability is a nearly linear function of the grain size, but that a large grain size requires a high degree of internal cleanness at the same time. The scatter is believed to be due to the different size and distribution of the nonmetallic inclusions, and to the additional influence of impurities, which are known to segregate at the grain boundaries and thus to hinder the grain growth, the most notorious of them being sulfur. Sulfur is objectionable anyway since, if not present segregated at grain boundaries, it will be present in the form of sulfides such as CaS or MgS giving rise to nonmetallic inclusions which again

Table 7-22 shows the nominal chemical compositions of some important commercial controlled expansion nickel-based alloys. Apart from alloy 2918, they are all essentially nickel-iron alloys with some manganese, silicon, and carbon as accompanying elements and, in some cases, additions of chromium and/or titanium combined with aluminum or niobium. Figure 7-33 shows the coefficient of linear expansion at 20 °C versus nickel content for nickel-iron alloys containing 0.4% Mn and 0.1% C (Wenschhof, 1980). The chromium additions are primarily for practical reasons, making it possible to melt the alloys in furnaces used for chromium-nickel steels, where the melt will take up chromium from the furnace lining. The additions of titanium combined with aluminum or niobium result in an increase in the hardness of the alloys on using an age-hardening treatment. Smaller amounts of aluminum, e.g., about 0 . 1 0.2% may also be present, being alloyed for the purpose of deoxidation during melting.

40

60 Mass % Ni

80

100

Figure 7-33. Coefficient of linear expansion at 20 °C vs. nickel content for nickel-iron alloys containing 0.4% Mn and 0.1% C (Wenschhof, 1980).

393

7.2 Composition, Structure, Properties, and Behavior

Table 7-22. Nominal chemical composition of some important commercial controlled expansion nickel-based alloys. Designation

Pernifer 36 40 42 4205 Ti 4206 42 Ti Nb 48 4706 50 51 5101 2918

Main alloying constituents (mass%) Ni

Fe

Mn

Si

C

36 40.6 42 40.6 42 43 48.3 47.3 50.8 51.5 51.3 29

63 59 57 48 52 52 51 46 48 48 47 54

0.25 0.5 0.5 0,4 0.2 0.1 0.4 0.2 0.5 0.2 0.5 0.2

0.15 0.15 0.15 0.6 0.2 0.1 0.15 0.15 0.20 0.05 0.10 0.1

0.005 0.005 0.005 0.03 0.02 0.01 0.015 0.01 0.005 0.005 0.005 0.005

The alloy in the first line of Table 7-22 has a coefficient of thermal expansion that is so low that its length is almost invariable for ordinary temperature changes around the ambient. For this reason the alloy was named "Invar". This behavior is caused by the magnetostriction which occurs during spontaneous magnetization below the Curie temperature which, for this alloy type, is about 200 °C (see Fig. 7-25). During cooling the increase in volume due to magnetostriction compensates for the decrease in volume due to thermal contraction until magnetic saturation is achieved, this effect being greatest around 0°C (Volk, 1970). The Invar alloys are also addressed in Vol. 3 of this Series. Heat-treatment and cold work change the expansivity of Invar considerably, as shown in Table 7-23. The expansivity is greatest for wellannealed material and least for quenched material. Cold drawing also decreases the expansivity. By cold working after quenching it is possible to produce material with a zero or even a negative coefficient of expansion. However, this behavior will not

Other

2.2 Ti 0.5 Al 5.3 Cr 5.6 Cr 2 Ti 0.5 Nb 2.3 Co 6Cr

1 Cr 17 Co

be stable over longer periods of time unless a stabilizing anneal has been performed, e.g., 2-4 hours at a temperature of at least 50-100 °C above the temperature during service, followed by slow cooling (Wenschhof, 1980). Substitution of 5 % cobalt for 5 % nickel provides an alloy with an expansion coefficient even lower than that of Invar. This

Table 7-23. Effect of heat-treatment on the coefficient of thermal expansion of Invara. Condition

Temperature (°Q

Mean coefficient (um/mK)

As-forged

17 to 100 17 to 250 18 to 100 18 to 250 16 to 100 16 to 250 16 to 100 16 to 250

1.66 3.11 0.64 2.53 1.02 2.43 2.01 2.89

Quenched from 830 °C (156O°F) Quenched from 830 °C and tempered Cooled from 830 °C to room temperature in 19 h Wenschhof(1980).

394

7 Nickel-Based Alloys

alloy, the so-called Superinvar, is also less susceptible to variations in the heat-treatment (Wenschhof, 1980). The Invar-type alloys, e.g., Pernifer 36, find applications wherever a very low coefficient of thermal expansion is required at not too high temperatures, e.g., compensating pendulums and balance wheels for clocks and watches, bimetal strip, vessels and piping for storage and transportation of liquefied natural gas, components for radios, TV, and other electronic devices, etc. The alloys of Table 7-22 with higher nickel contents have higher coefficients of thermal expansion in ordinary earth environments, as shown in Fig. 7-34. These alloys including 2918 have coefficients of thermal expansion close to those of glasses and ceramics. They therefore find uses as so-called glass seal-in wires in connection with soft glasses (alloys with 47-54% nickel and alloy 4206) or hard glasses (alloy 2918) and ceramics (alloy 2918 and alloys with up to 43 % nickel). Typical applications include integrated-circuit lead

12.0

100

200

500 300 400 Temperature I C)

600

700

Figure 7-34. Average linear coefficient of thermal expansion of various nickel-based controlled expansion alloys between 20 °C and the temperature indicated in the diagram (data from Krupp VDM GmbH).

frames, reed relays, incandescent lamps, and vacuum tubes. Alloy 2918 is also used in aerospace applications, where a martensitic transformation has to be avoided even at very low temperatures of down to —196 °C. Figure 7-34 shows the average linear coefficient of thermal expansion of this alloy group between 20 °C and the temperature indicated in the diagram. According to the actual requirements of the electronics industry, additional alloy variations can be created, having different nickel or other alloying element contents according to their influence on the expansivity (Wenschhof, 1980).

7.3 Concluding Remarks The applications of nickel-based alloys as mentioned in this chapter will continue in the future. In the field of wet corrosive applications there is still ample room for new alloy developments triggered by new requirements for materials as a result of more stringent pollution control legislation and new and more efficient chemical processes. Where in this field nickel-based alloys encounter competition from plastic and resin-based materials, they offer the advantages of longer life and less problems with respect to recycling. Engineering ceramics may well constitute an addition to today's materials world for service at extremely high temperatures in the distant future, but nickel-aluminides (see Chap. 11 of this Volume) may be the new materials for high-temperature service in the short-term especially where their excellent resistance to the corrosive attack of various gaseous media (Brill and Klower, 1991) is a primary requirement. While nickel-based and iron-based oxide-dispersion-strengthened superalloys have entered the market during the last decade

7.4 References

with application potentials for up to 1150°C or even 1350°C (Riihle and Korb, 1991), their actual shortcomings with respect to size limitations, high cost, weldability, and low-cycle fatigue impose severe restrictions upon their use. Therefore the nickel-based high-temperature alloys described in this Chapter and their future derivatives will remain the workhorses for service at intermediate and high temperatures of up to 950 or even 1200°C, depending on the requirements of creep strength and service life. Along with dispersion strengthening and intermetallic phases, the nickel-based high-temperature alloys still offer considerable potential for future developments (Arzt, 1991). As a result of the increasing use of electronics, the soft magnetic nickel-iron alloys will gain more importance in the future, and hence new alloy developments will be seen in this field. Here the amorphous and nanocrystalline metals constitute an addition to the classic world of metallic materials offering the advantages of lack of magnetocrystalline anisotropy and plastic deformability. However, their limitations with respect to cost of manufacturing and size will keep the broader fields of application open to the crystalline nickel-iron alloys.

7.4 References Agarwal, D. C , Brill, U. (1993 a), in: Corrosion 93. Houston, TX: NACE International, Paper No. 226. Agarwal, D. C , Brill, U. (1993 b), in: Corrosion 93. Houston, TX: NACE International, Paper No. 209. Agarwal, D. C , Heubner, U., Kirchheiner, R., Koehler, M. (1991), in: Corrosion 91. Houston, TX: National Association of Corrosion Engineers, Paper No. 179. Antolovich, S. D., Stusrud, R. W., MacKay, R. A., Anton, D. L., Khan, T, Kissinger, R. D., Klarstrom, D. L. (Eds.) (1992), Super alloys 1992, Proc. 7th Int. Symp. on Superalloys. Warrendale, PA: TMS.

395

Arzt, E. (1991), in: Symp. Materialforschung des BMFT. Cologne: TUV Rheinland, pp. 102-116. Betteridge, W, Heslop, I (Eds.) (1974), The Nimonic Alloys, 2nd Ed. London: Edward Arnold. Brill, U. (1990), Werkst. Korr. 41, 682-688. Brill, U. (1992), Metall 46, 778-782. Brill, U. (1992), Korrosion von Nickel, Cobalt und Nickel- und Cobalt-Basislegierungen, preprint from Korrosion und Korrosionsschutz, 2nd ed.: Kunze, E. (Ed.). Berlin: Walter de Gruyter. Brill, U. (1995), in: Nickel Alloys and High Alloy Special Stainless Steels, 2nd ed: U. Heubner (Ed.), Ehningen: expert-Verlag Brill, U., Agarwal, D. C. (1993), in: Corrosion 93. Houston, TX: NACE International, Paper No. 226. Brill, U., Klower, J. (1991), in: Mater. Res. Soc. Symp. Proc, Vol. 213: L. A. Johnson, D. P. Pope, J. O. Stiegler (Eds.), Pittsburgh, PA: MRS, pp. 963-968. Brill, U., Klower, J. (1993), Mater. High Temp. 11, 151-158. Brill. U., Heubner, U., Drefahl, K., Henrich, H. J. (1991a), Ingenieur-Werkstoffe 3, 59-62. Brill, U., Heubner, U., Rockel, M. (1991b), Metall 44, 936-946. Bryant, I R., Wallis, E. (1988), in: Corrosion 88. Houston, TX: The National Association of Corrosion Engineers, Paper No. 68. Bryant, J. R., Koehler, M., Heubner, U., Rice, P. (1990), in: Corrosion 90. Houston, TX: The National Association of Corrosion Engineers, Paper No. 51. Buth, X, Kohler, M. (1992), Werkst. Korr. 43, 421425. Coppolecchia, V., Bryant, X, Hofmann, R, Drefahl, K. (1987), in: Performance of High Temperature Materials in Fluidized Bed Combustion Systems and Process Industries: Ganesan, P., Bradley, R. A. (Eds.), Materials Park, OH: ASM Int., pp. 201209. Drefahl, K., Hofmann, F. (1986), Paper presented at ASM Materials Week 1986, Orlando, FL. Drefahl, K., Matucha, K. H., Hofmann, F. (1986), in: Corrosion 86. Houston, TX: The National Association of Corrosion Engineers, Paper No. 376. Duhl, D. N., Maurer, G., Antolovich, S., Lund, C , Reichman, S. (Eds.) (1988), Superalloys 1988, Proc. 6th Int. Symp. on Superalloys. Warrendale, PA: TMS. Frank, R. B., DeBold, T. A. (1988), Mater. Performance 27 (9), 59-66. Friend, W. Z. (1980), Corrosion of Nickel and NickelBase Alloys. New York: Wiley. Ganesan, P., Clatworthy, E. R, Harris, X A. (1987), in: Corrosion 87. Houston, TX: National Association of Corrosion Engineers, Paper No. 286. Gell, M., Kortovich, C. S., Bricknell, R. H., Kent, W. B., Radavich, X R (Eds.) (1984), Superalloys 1984, Proc. 5th Int. Conf on Superalloys. Warrendale, PA: TMS-AIME.

396

7 Nickel-Based Alloys

Hattendorf, H. (1991), unpublished, Krupp VDM GmbH, Werdohl, Germany Heubner, U. (Ed.) (1987 a), Nickel Alloys and HighAlloy Special Stainless Steels. Sindelfingen, Germany: expert-Verlag. Heubner, U. (1987 b), in: Handbuch der Fertigungstechnik, Bd. 4\2 Wdrmebehandeln: Spur G. (Ed.). Munich: Carl Hanser Verlag, pp. 972-1011. Heubner, U. (Ed.) (1995), Nickel Alloys and High Alloy Special Stainless Steels, 2nd ed. Ehningen, Germany: expert-Verlag. Heubner, U., Hofmann, F. (1989), Werkst. Korr. 40, 363-369. Heubner, U., Kirchheiner, R. (1987), in: Nickel Alloys and High-Alloy Special Stainless Steels: Heubner, U. (Ed.). Sindelfingen, Germany: expert-Verlag, pp. 223-240. Heubner, U., Kohler, M. (1992), Werkst. Korr. 43, 181-190. Heubner, U., Kohler, M. (1994), in: Superalloys 718, 625, 706 and Various Derivates, Proc. Int. Symp.: Loria, E. A. (Ed.). Warrendale, PA: TMS, pp. 479488. Heubner, U., Rockel, M. (1986), Werkst. Korr. 37, 7-12. Heubner, U., Rockel, M. B., Wallis, E. (1985), ATB Metallurgie 25, 235-241. Heubner, U., Heimann, W., Kirchheiner, R. (1987), Werkst. Korr. 38, 746-156. Heubner, U., Rockel, M., Wallis, E. (1989 a), Werkst. Korr. 40, 418-426. Heubner, U., Rockel, M., Wallis, E. (1989 b), Werkst. Korr. 40, 459-466. Heubner, U. L., Altpeter, E., Rockel, M. B., Wallis, E. (1989c), Corrosion-NACE 45, 249-259. Heubner, U., Hoffmann, T., Rudolph, G. (1990), in: Weldability of Materials: Patterson, R. A., Mahin, K. W. (Eds.). Materials Park, OH: ASM Int., pp. 175-182. Heubner, IL, Drefahl, K., Henrich, H. J. (1991 a), in: Proc. 1st Int. Conf. on Heat Resistant Materials: Natesan, K., Tillack, D. J. (Eds.). Materials Park, OH: ASM Int., pp. 495-504. Heubner, U., Kirchheiner, R., Rockel, M. (1991b), in: Corrosion 91. Houston, TX: The National Association of Corrosion Engineers, Paper No. 321. Hibner, E. L. (1991), in: Corrosion 91. Houston, TX: National Association of Corrosion Engineers, Paper No. 18. Holt, R. T, Wallace, W. (1976), Int. Met. Rev. 203. Kirchheiner, R., Stenner, F. (1993), in: Corrosion 93. Houston, TX: NACE Int., Paper No. 417. Kirchheiner, R., Schalk, W, Muller, H., Palomino, S. M. (1989), Werkst. Korr. 40, 545-551. Kirchheiner, R., Koehler, M., Heubner, U. (1990), in: Corrosion 90. Houston, TX: The National Association of Corrosion Engineers, Paper No. 90. Kirchheiner, R., Kohler, M., Heubner, U. (1992), Werkst. Korr. 43, 388-395.

Klarstrom, D. L., Sridhar, N., Kolts, X, Kiser, S. D., Crum, J. R. (1987), in: Metals Handbook, 9th ed. Vol. 13, Corrosion: Davis, J. R., Destefani, J. D., Frisell, H. J., Crankovic, G. M. (Eds.). Metals Park, OH: ASM Int., pp. 641-657. Kofstad, P. (1988), High Temperature Corrosion. London: Elsevier, pp. 401-407; 514-516; 534. Kohler, M. (1991), in: Superalloys 718, 625 and Various Derivates, Proc. Int. Symp., June 1991: Loria, E. A. (Ed.). Warrendale, PA: TMS, pp. 363-374. Kohler, M., Heubner, U. (1994), in: Corrosion 94. Houston, TX: NACE Int., Paper No. 230. Krause, H. H. (1988), Corrosion 88. Houston, TX: The National Association of Corrosion Engineers, Paper No. 5. Krupp VDM GmbH (Ed.) (1992), High-Performance Materials. VDM-Publication N 524 92-11. Werdohl, Germany: Krupp VDM. Lee III, T. S., Hodge, F. G. (1976), Mater. Performance 9 (15), 29-36. Loria, E. (Ed.) (1989), Superalloy 718, Metallurgy and Application, Proc. Int. Symp. June 1989. Warrendale, PA: TMS. Loria, E. (Ed.) (1991), Superalloys 718, 625 and Various Derivatives, Proc. Int. Symp. Warrendale, PA: TMS. Loria, E. (Ed.) (1994), Superalloys 718, 625, 706 and Various Derivatives, Proc. Int. Symp. Warrendale, PA: TMS. McLean, ML, Strang, A. (1984), Met. TechnoL 11, 454-464. Massalski, T. B. (1991), Binary Alloy Phase Diagrams, 2nd ed. Materials Park, OH: ASM Int. Morinaga, M., Yukawa, N., Adachi, H., Ezaki, H. (1984), in: Superalloys 1984: Gell, M., Kortovich, C. S. Bricknell, R. H., Kent, W. B., Radavich, J. F. (Eds.). Metals Park, OH: ASM. Muzyka, D. R. (1972), in: The Superalloys: Sims, C. T, Hagel, W. C. (Eds.). New York, Wiley, pp. 113-143. Pfeifer, E, Radeloff, C. (1980), J. Magn. Magn. Mater. 19, 190-207. Pfeiffer, H., Thomas, H. (1963), Zunderfeste Legierungen, Berlin: Springer, pp. 275-283. Raghavan, M., Mueller, R. R., Vaughn, G. A., Floreen, S. (1984), Met. Trans. 15A, 783-792. Renner, M., Heubner, U., Rockel, M. B., Wallis, E. (1986), Werkst. Korr. 37, 183-190. Rivlin, V. G., Raynor, G. V (1980), Int. Met. Rev. 1, 21-38. Rockel, M. (1987), in: Nickel Alloys and High-Alloy Special Stainless Steels: Heubner, U. (Ed.). Sindelfingen, Germany: expert-Verlag, pp. 48-77. Rosa, E., Smith, G. D. (1987), in: Performance of High Temperature Materials in Fludized Bed Combustion Systems and Process Industries: Ganesan, P., Bradley, R. A. (Eds.). Materials Park, OH: ASM Int., pp. 149-153.

7.4 References

Riihle, M., Korb, G. (1991), in: Symp. Materialforschung des BMFT. Cologne: TUV Rheinland, pp. 702-727; in: Proc. 1st Int. Conf. on Heat Resistant Materials: Natesan, K., Tillack, D. J. (Eds.). Materials Park, OH: ASM Int., pp. 45-59. Sedriks, A. J. (1979), Corrosion of Stainless Steels. New York: Wiley, pp. 20, 172-178, 210-231, 232240. Sims, C. T., Hagel, W. C. (Eds.) (1972), The Superalloys. New York: Wiley. Sims, C. T., Stoloff, N. S., Hagel, W. C. (Eds.) (1987), The Superalloys II. New York: Wiley. Taylor, A. (1956), Trans. AIME 206, 1356-1362. VDM Nickel-Technologie AG (Ed.) (1990), VDMProdukte fur den Industrieofenbau, Publication N 501. Werdohl, Germany: Krupp VDM. Volk, K. E. (Ed.) (1970), Nickel und Nickellegierungen. Berlin: Springer. Wenschhof, D. (Ed.) (1980), Metals Handbook, 9th Ed., Vol. 3. Metals Park, OH: ASM Int., pp. 792798. Woodyatt, L. R., Sims, C. T., Beattie, H. J. (1966), Trans. AIME 236, 519-527.

397

General Reading Agarwal, D. C , Heubner, U., Kohler, M., Herda, W. (1994), Mater. Perform. 33 (10), 64-68. Bradley, E. F. (1988), Superalloys - A Technical Guide. Metals Park, OH: ASM Int. Friend, W. Z. (1980), Corrosion of Nickel and NickelBase Alloys. New York: Wiley. Heubner, U. (Ed.) (1965), Nickel Alloys and High Alloy Special Stainless Steels, 2nd ed. Ehningen, Germany: expert-Verlag. Heubner, U., Kohler, M. (1992), Werkst. Korr. 43, 181-190. Kirchheiner, R., Kohler, M., Heubner, U. (1992), Werkst. Korr. 43, 388-395. Kofstad, P. (1988), High Temperature Corrosion. London: Elsevier. Natesan, K., Tillack, D. J. (Eds.) (1991), Heat Resistant Materials, Proc. 1st Int. Conf. on Heat-Resistant Materials. Materials Park, OH: ASM Int. Sims, C. T, Hagel, W C. (Eds.) (1972), The Superalloys. New York: Wiley. Sims, C. T, Stoloff, N. S., Hagel, W C. (Eds.) (1987), The Superalloys II. New York: Wiley.

8 Titanium, Zirconium, and Hafnium Francis H. (Sam) Froes Institute for Materials and Advanced Processes, University of Idaho, Moscow, ID, U.S.A. Te-Lin Yau Teledyne Wah Chang, Albany, OR, U.S.A. Hans G. Weidinger Nuremberg, Germany

List of 8.1 8.2 8.2.1 8.2.2 8.2.2.1 8.2.2.2 8.2.2.3 8.2.2.4 8.2.2.5 8.2.3 8.2.3.1 8.2.3.2 8.2.3.3 8.2.3.4 8.2.3.5 8.2.3.6 8.2.3.7 8.2.3.8 8.2.3.9 8.2.4 8.2.4.1 8.2.4.2 8.2.4.3 8.2.4.4 8.2.4.5 8.2.4.6 8.2.4.7

Symbols and Abbreviations Introduction Titanium History General Characteristics Physical Properties Alloying Behavior Phase Diagrams Alloy Classes Microstructural Development Processing/Fabrication Extraction Ingot Castings Powder Metallurgy Joining Wrought Product Processing Wrought Product Forming Machining Metal Matrix Composites Mechanical Properties Cast and Wrought Terminal Alloys Intermetallic Alloys Cast Alloys Powder Metallurgy Alloys Welded Components Machined Components Metal Matrix Composites

Materials Science and Technology Copyright © WILEY-VCH Verlag GmbH & Co KGaA. Allrightsreserved.

401 403 403 403 405 405 405 407 408 409 412 412 413 413 413 414 414 415 415 416 417 417 421 423 426 428 428 428

400

8.2.5 8.2.6 8.2.7 8.3 8.3.1 8.3.2 8.3.2.1 8.3.2.2 8.3.2.3 8.3.2.4 8.3.3 8.3.4 8.3.5 8.3.6 8.3.6.1 8.3.6.2 8.3.7 8.4 8.4.1 8.4.2 8.4.2.1 8.4.2.2 8.4.3 8.4.4 8.4.5 8.4.6 8.4.6.1 8.4.6.2 8.4.7 8.5

8 Titanium, Zirconium, and Hafnium

Chemical Properties/Corrosion Behavior Applications Concluding Remarks and Future Thoughts Zirconium History General Characteristics Physical Properties Alloy Behavior Phase Diagrams Microstructural Development Processing/Fabrication Mechanical Properties Chemical Properties/Corrosion Behavior Applications Nuclear Industrial Concluding Remarks and Future Thoughts Hafnium History General Characteristics Physical Properties Alloy Behavior Processing/Fabrication Mechanical Properties Chemical Properties/Corrosion Behavior Applications Nuclear Industrial Concluding Remarks and Future Thoughts References

428 431 434 436 436 438 439 439 440 441 443 443 444 447 447 452 458 459 459 460 460 460 460 461 461 463 463 464 466 466

List of Symbols and Abbreviations

List of Symbols and Abbreviations A

Kt N n s/N t

constant relative atomic mass grain size after annealing initial grain size fracture toughness notch concentration number of cycles constant stress/number of cycles to failure time

ASME b.c.c. BE BWR CHIP CPI DB DM FA FCGR FL HCF h.c.p. HIP'ing HSSCC IM ksi LCF LWR MA MIBK MMC NG NL PA PCI PM PM/RS PWR RA RFNA RMBK

American Society of Mechanical Engineers body-centered cubic blended elemental boiling water reactor cold and hot isostatic pressing chemical process industry diffusion bonding duplex fuel assembly fatigue crack growth rate fully lamellar high cycle fatigue hexagonal close packed hot isostatic pressing hot salt stress-corrosion cracking ingot metallurgy kilopounds per square inch low cycle fatigue light water reactor mechanical alloying methyl isobutyl ketone metal matrix composites near gamma near lamellar pre-alloyed pellet-cladding interaction powder metallurgy powder metallurgy/rapid solidification pressurized water reactor reduction in area red fuming nitric acid Russian-type water cooled reactor

A D Do

401

402

RS SCC SPF SS ST TCP TIG UNS UTS VD VVER YS

8 Titanium, Zirconium, and Hafnium

rapidly solidified stress corrosion cracking superplastic forming stainless steel solution treatment thermochemical processing tungsten inert gas Unified Numbering System ultimate tensile strength vapor deposition Russian-type water cooled reactor yield strength

8.2 Titanium

8,1 Introduction This chapter deals with the three metals titanium, zirconium and hafnium. While all three metals have many similarities because of their position in the periodic table, they also show very significant differences. All three metals exhibit excellent corrosion behavior leading to chemical processing industry applications. Because of its low density and very attractive fracture related properties, titanium is used extensively in aerospace structural applications. In contrast, zirconium and hafnium are not low density materials but exhibit complementary properties for nuclear applications: zirconium is transparent to and hafnium is opaque to thermal neutrons. Structural applications require not only a precise chemistry, but also careful control of microstructure, and hence processing route, to optimize mechanical property combinations. Thus in the following sections, the reader will note a significant emphasis on topics such as processing, microstructure, mechanical properties in the titanium portion of this chapter, and much less on these issues in the case of the zirconium and hafnium.

8.2 Titanium 8.2.1 History Titanium has been recognized as an element for 200 years. However, it is only in the last 45 years that it has gained strategic importance. During that time, commercial production of titanium and titanium alloys in the United States has increased from zero to a peak of 25 million kg/yr. The catalyst for this remarkable growth was the development by Dr. Wilhelm X Kroll of a relatively safe, economical

403

method to produce titanium metal in the late 1930s. KrolPs process involved reduction of titanium tetrachloride (TiCl4), first with sodium and calcium, and later with magnesium, under an inert gas atmosphere (Bomberger et al., 1985). Research by Kroll and many others continued through World War II. By the late 1940s, the mechanical properties, physical properties, and alloying characteristics of titanium were defined as the commercial importance of the metal became apparent. In 1949, the first titanium for actual flight purposes was ordered from Remington Arms (later RemCru, and still later, Crucible Steel) in the U.S.A. by Douglas Aircraft. Other early U.S.A. entries to the titanium field were made by Mallory-Sharon (later RMI) and TMCA (later Timet). In the U.K., ICI Metals (later IMI) began sponge production in 1948, with further involvement from continental Europe a few years later. Recognizing the military potential of titanium, the Soviets began sponge production in 1954. In Japan, sponge production was initiated by Osaka Titanium in 1952, generally with the aim to supply other countries. In the U.S.A. the government invested large amounts of money to develop the science and technology of titanium and its alloys. This included an investment of more than a quarter billion dollars for a sponge stock-pile up to 1964, and the establishment of a titanium laboratory at Battelle in 1955 (Bomberger et al., 1985). The vast influx of money resulted in the rapid development of a sound technology with very good scientific underpinning. The double consumable arc melting technique was developed by Armor Research Foundation in 1953. Problems relating to hydrogen embrittlement and hot salt stress corrosion cracking were recognized and circumvented.

404

8 Titanium, Zirconium, and Hafnium

Alloy development progressed rapidly from about 1948 with people at Remington Arms recognizing the beneficial effects of aluminum additions. The "work-horse" Ti-6A1-4V 1 alloy was introduced in 1954, with the disputed patent assigned to Rem- Cru. This alloy soon became by far the most important titanium alloy because of its excellent combination of mechanical properties and "forgiving" processability. The first p titanium alloy (Ti-13 V - l 1 C r 3A1) was developed by Rem-Cru in the 1950s, with this high strength heat treatable alloy seeing extensive use in the SR71 high speed surveillance aircraft. Alloy development in the U.K., driven by RollsRoyce, was concentrated more on elevated temperature alloys for use in engines. Aircraft manufacturers have employed generally increasing amounts of titanium in airframe applications for heavily stressed demanding components. In the U.S.A. engine components using titanium first appeared with the Pratt and Whitney J57 in 1954, including discs, blades and spacers in the compressor sector; in the U.K. it was the Rolls-Royce Avon engine in 1954. Work in the 1950s also indicated the excellent corrosion resistance of titanium and its alloys, and early commercial applications included use in anodizing, wet chlorine, and nitric acid equipment. During the late 1950s the shipment of titanium mill products in the U.S.A. dropped from a high of 4.5 million kg in 1957 due to a change in emphasis from aircraft to missiles, and concerns over the cost of titanium. Mill shipments in the U.S.A. tripled during the 1960s from 4.5 million kg in 1960, with aerospace use accounting for 90 % of the market in 1970. This increase resulted 1 Throughout the text, all terminal alloys are given in wt. %, intermetallics in at. %.

mainly from use in non-military engines and wide-body jets, such as the Boeing 747, the DC-10 and the L-1011. Advances also occurred in melting practices and expanded use in non-aerospace markets such as ships and heat-exchanger tubing. The 1971 cancellation of the Mach 3 Supersonic Transport (SST), which was slated to use considerable amounts of titanium, was a blow to the titanium industry with mill products in the U.S.A dropping from 12 million kg in 1970 to 9 million kg in 1971. Mill product shipments increased in the 1970s for the most part as a result of increased use in the large commercial transports and the necessary high bypass engines, new military airframes with 20-35% of their structural weight produced from titanium products, and nonaerospace use, a result of the excellent corrosion resistance of titanium. The U.S. industry set a new record in 1980-1981 of about 23 million kg, but this figure then dropped because of hedge buying by the aerospace industry. This cyclic nature of the titanium industry will only smooth out if non-aerospace use increases. In Europe, and to a greater extent in Japan, industrial applications exceed 50% of the total use. Titanium mill shipments in the U.S.A. increased steadily during the Reagan era, and early years of the Bush administration with the build up in military hardware, peaking in 1990 at a record 25 million kg. When "peace broke out", symbolized by the dismantling of the Berlin Wall, titanium shipments fell precipitously soon thereafter to about 16 million kg per year; a level which may well be the "norm". This has left a great over-capacity in the U.S.A., and a painful "tailoring" by the titanium industry. Now the expansion of the titanium market will be even more dependent on reducing the cost; as a consequence

8.2 Titanium

lower cost, albeit inferior performance, alloys are already being introduced into the marketplace. A complicating factor is going to be the effect of generally low-cost products from the former U.S.S.R.; where the peak capacity is estimated to have been four times that of the U.S.A., i.e., as much as 90 million kg of mill products per year. 8.2.2 General Characteristics

Titanium alloys may be divided into two major categories: corrosion resistant and structural alloys. The corrosion resistant alloys are generally based on a single phase a with dilute additions of solid solution strengthening and a-stabilizing elements like oxygen, palladium, and aluminum. These alloys are used in the chemical, energy, paper and food processing industries to produce highly corrosion resistant tubings, heat exchangers, valve housings, and containers. In addition to excellent corrosion resistance, the single phase alpha alloys provide good weldability, and easy processing and fabrication, but relatively low strength. The structural alloys can be divided into four categories: the near a alloys, the a + /? alloys, the /? alloys and the titanium aluminide intermetallics, which will be discussed later in Sec. 8.2.2.4. In markets served by major U.S. titanium producers, corrosion resistant alloys make up 25% of the total output, T i 6A1-4V, 60% and all other structural alloys, the remaining 15%. The attractive combinations of mechanical properties which can be obtained in titanium alloys make them prime candidates for many aerospace and commercial applications. However, titanium alloy components are expensive, a fact that often limits use. For example, Table 8-1 gives the amount of titanium slated for use in

405

Table 8-1. Airframe weight percentage of titanium. System C5 (cargo) Bl (bomber) F15 (fighter)

Early design

Final concept

24 42 50

3 22 34

three U.S. Air Force systems, expressed as airframe weight percentage, with early design figures shown for comparison (Bomberger et al., 1985). Thus, much work has been focused on reducing component cost while maintaining acceptable mechanical property levels as well as on approaches that include near net shape techniques and lower cost alloy formulation. 8.2.2.1 Physical Properties Titanium in its natural form is a dark gray color; however it is easily anodized to give a very attractive array of colors leading to use in jewelry, for decoration of buildings, and for various other applications where appearance is important. The metal and its alloys exhibit low density (approximately 60% of the density of steel). Titanium is nonmagnetic and has good heat-transfer characteristics. Its coefficient of thermal expansion is somewhat lower than that of steel, and less than half that of aluminum. The melting points of titanium and its alloys are higher than that of steel, but the maximum use temperature is much lower than would be anticipated based on this characteristic alone. A summary of the physical (and a few mechanical properties) is given in Table 8-2. 8.2.2.2 Alloying Behavior Titanium exists in two crystalline states : in a low-temperature alpha (a) phase which has a close-packed hexagonal crystal structure, and a high temperature beta (/?) phase

406

8 Titanium, Zirconium, and Hafnium

Table 8-2. Physical and mechanical properties of elemental titanium. Atomic number Atomic weight Atomic volume Covalent radius First ionization energy Thermal neutron absorption cross section Crystal structure

Color Density Melting point Solidus/liquidus Boiling point Specific heat (at 298 K) Thermal conductivity Heat of fusion Heat of vaporization Specific gravity Hardness Tensile strength Modulus of elasticity Young's modulus of elasticity Poisson's ratio Coefficient of friction Specific resistance Coefficient of thermal expansion Electrical conductivity Electrical resistivity Electronegativity Temperature coefficient of electrical resistance Magnetic susceptibility Machinability rating

22 47.90 10.6 weight/density 0.132 nm 661.5 MJ/(kgmol) 560 fm2/atom • a: close-packed, hexagonal < 1156 K • p: body-centered, cubic >1156K Dark gray 4510 kg/m3 1941 ±285 K 1998 K 3533 K 0.518 J/(kgK) 21 W/(m K) 440 kJ/kg 9.83 MJ/kg 4.5 HRB 70 to 74 241 GPa 102.7 GPa 102.7 GPa 0.41 0.8 at 40 m/min 0.68 at 300 m/min 0.554 uQ m 8.64 x l O ^ K " 1 3% IACS (copper 100%) 0.478 uQ m 1.5 Pauling's 0.0026 K" 1 1.25xlO~ 6 emu/g 40 (equivalent to 3/4 hardness stainless steel)

which has a body centered cubic structure (Fig. 8-1) (Joseph and Froes, 1988). The allotropic transformation occurs at 882 °C in nominally pure titanium. Titanium has certain features which make it very different from other light metals such as aluminum and magnesium (Polmear, 1989).

Body-centered cubic

Beta Transus Temperature 882 °C

Hexagonal close packed

Figure 8-1. The two allotropic forms of titanium. The transition from the low temperature a phase to the high temperature ($ phase occurs at 882 °C (Joseph and Froes, 1988).

The allotropic transformation allows for formation of alloys composed of a, jS, or a/j8 microstructures in addition to compound formation. Because of its electronic structure, titanium can form solid solutions with most substitutional elements having a size factor within 20%, giving the opportunity for many alloying possibilities. Titanium also reacts with interstitial elements such as nitrogen, oxygen, and hydrogen. These reactions may occur at temperatures below the melting point. When reacting with other elements, titanium may form solid solutions and compounds with metallic, covalent or ionic bonding. The choice of alloying elements is determined by the ability of the element to stabilize either the a- or j8-phases (Fig. 8-2) (Molchanova, 1965). Aluminum, oxygen, nitrogen, gallium, and carbon are the most common a-stabilizing elements. Zirconium, tin, and silicon are regarded as neutral in their ability to stabilize either phase. Elements which stabilize the j8 phase can either form binary systems of the j8-iso-

8.2 Titanium

407

Binary Titanium Alloys

a-stabilized

^-stabilized

Eutectoid Trans formation (^•eutectoid)

Simple Peritectic

Solutes

Solutes

Solutes

V Zr Nb Mo Hf Ta Re

Cr Mn Fe Co Ni Cu Pd Ag W Pt Au

Simple Transformation (0-isomorphous)

Peritectoid (5-»c£ Transformation (&~peritectoid)

Solutes B,Sc, Ga, La Ce, Gd, Nd, Ge Al, C

N, 0

H, Be,Si,Sn, Pb, Bi, U

ot + y

OL + y

77

Ti

Ti Solute Content

»*»

Figure 8-2. Classification scheme for binary titanium alloys (Molchanova, 1965; reproduced with permission from The Physical Metallurgy of Titanium Alloys, ASM Int., Materials Park, OH 1984, p. 43).

morphous type or the jS-eutectoid type (see next section). Elements forming the isomorphous type binary system include Mo, V, and Ta, while Cu, Mn, Cr, Fe, Ni, Co, and H are eutectoid formers. The /?-isomorphous alloying elements, which do not form intermetallic compounds, have traditionally been preferably added to the eutectoid-type elements as addition to a-p or p alloys to improve hardenability and increase response to heat treatment.

8.2.2.3 Phase Diagrams A number of research groups have attempted to categorize titanium alloy phase diagrams (Margolin and Neilson, 1960; Molchanova, 1965), all agreeing on two major subdivisions: a-stabilized and /?-stabilized systems. Of these perhaps the most convenient is that developed by Molchanova (Molchanova, 1965), Fig. 8-2. Here the a stabilizers are divided into those having perfect stability, in which the a phase can coexist with the liquid (e.g., T i - O and Ti-N) and there is a simple peritectic reac-

408

8 Titanium, Zirconium, and Hafnium

tion, and those which have limited a stability in which, with decreasing temperature, decomposition of the a takes place by a peritectoid reaction into p phase plus a compound (P peritectoid). Examples of the latter type of system are Ti-B, Ti-C, and Ti-Al. Molchanova also divides the P stabilizers into two categories, p isomorphous and /? eutectoid. In the former system an extreme P solubility range co-exists with only restricted a solubility. Examples are Ti-Mo, Ti-Ta, Ti-V, with elements such as Zr and Hf occupying an intermediate position since they have complete mutual solubility in both the a and /? phases. For the P eutectoid systems the p phase has a limited solubility range and decomposes into the a phase and a compound (e.g., Ti-Cr and Ti-Cu). This class can also be further subdivided depending on whether the P transformation is rapid (the "active" eutectoid formers such as Ti-Si, Ti-Cu, and Ti-Ni) or slow (the "sluggish" eutectoid formers such as Ti-Cr and Ti-Fe). 8.2.2.4 Alloy Classes Alloy Categories Titanium alloys are categorized into four groups: a, a-/?, P alloys, and intermetallics (77XA1, where x = l or 3). Titanium alloys for aerospace applications contain a- and /^-stabilizing elements to achieve the required mechanical properties such as tensile strength, creep, fatigue, fatigue crack propagation resistance, fracture toughness, stress-corrosion cracking, and resistance to oxidation (Froes et al, 1985). Once the chemistry is selected, optimization of mechanical properties is achieved by working to control the size, shape and dispersion of both the /? and a phases. /?-isomorphous alloying elements (e.g., Mo, V, Nb) which do not form intermetallic compounds have traditionally been pre-

ferred to "eutectoid-type" elements (e.g., Cr, Cu, Ni). However some /J-eutectoid type compound formers are added to a-fi or p alloys to improve hardenability and increase response to heat treatment. a Alloys The a alloys contain predominantly a phase at temperatures up to well above 540 °C (1000 °F). A major class of a alloys is the unalloyed-titanium family of alloys, which differ in the amount of oxygen and iron in each alloy. Alloys with higher interstitial content are higher in strength, hardness, and transformation temperature compared to high purity alloys. Other a alloys contain additions such as Al and Sn (e.g., Ti-5Al-2.5Sn and T i - 6 A l - 2 S n - 4 Z r 2Mo). Generally, a-rich alloys are more resistant to high temperature creep than a-P or P alloys, and a alloys exhibit little strengthening by heat treatment. These alloys are usually annealed or recrystallized to remove stress from cold working; they have good weldability and generally inferior forgeability in comparison to a-/? or P alloys. a-P Alloys (x-P alloys contain one or more of the a and P stabilizers. These alloys retain more p after final heat treatment than the near a alloys, and can be strengthened by solution treating and aging, although they are generally used in the annealed condition. Solution treatment is usually performed high in the a-p phase field followed by aging at lower temperatures to precipitate a, giving a mixture of fine a in a /? matrix. The solution treating and aging can increase the strength of these alloys by up to 80% (Froes et al., 1985). Alloys with low amounts of p stabilizer (for example T i -

8.2 Titanium

6A1-4V) have poor hardenability and must be rapidly quenched for subsequent strengthening. Water quenching of T i 6A1-4V will adequately harden sections of only less than 25 mm. p stabilizers in a-/? alloys increase hardenability. p Alloys p alloys have more P stabilizer content and less a stabilizer than a-/? alloys. These alloys have high hardenability with the P phase retained completely during air cooling of thin sections and water quenching of thick sections. P alloys have good forgeability and good cold formability in the solution-treated condition. After solution treatment, aging is performed to transform some j8 phase to a. The strength level of these alloys is greater than a-/? alloys, a result of the finely dispersed a particles in the p phase. These alloys have relatively higher densities, and generally lower creep strengths than the a-/? alloys. The fracture toughness of aged p alloys at a given strength level is generally higher than that of an aged a-p alloy, although crack growth rates can be faster. Titanium Aluminides To increase the efficiency of gas turbine engines, higher operating temperatures are necessary, requiring alloys with enhanced mechanical properties at elevated temperatures. The family of titanium alloys showing potential for applications at as high at 900 °C are the titanium aluminide intermetallic compounds Ti3Al(a2) and TiAl(y) (Froesetal. 1991; Froesetal., 1992;Froes, 1994; see also Chap. 11 by Sauthoff in this Volume). The major disadvantage of this alloy group is their low room temperature ductility. However, it has been found that niobium, or niobium with other p stabilizing elements, in combination with mi-

409

crostructure control, can increase room temperature ductility in the Ti3Al alloys up to as much as 26% elongation. Recently, by careful control of the microstructure the ambient temperature ductility of two-phase TiAl(y + a2) has been raised to almost 5 % elongation. However, major challenges remain, particularly with the TiAl compositions, relating to characteristics such as fracture toughness and fatigue crack growth rate. 8.2.2.5 Microstructural Development

In addition to chemistry, the mechanical properties of titanium alloys are strongly influenced by the microstructure (Froes et al., 1985). In turn, the microstructure is critically dependent on the processing, particularly whether this is carried out above or below the p transus temperature, the temperature below which the a phase is stable. In general terms two microstructural features are important in commercial terminal alloys: (a) the p grain size and shape, and (b) the morphology of the a phase within the P grains. Similar features strongly influence the properties of the intermetallics, but a discussion of these features is beyond the scope of this article; the interested reader is referred to Froes et al. (1991); Froes et al. (1992); Ward (1993); Kim (1994); Froes (1994). P Grains

Control of the P grain size is dependent on two factors, recrystallization (when this occurs because of sufficient working) and subsequent grain growth (Froes et al., 1985). A number of techniques have been developed for recrystallization of the p grains in a and a-/? alloys by working followed by high P field annealing. The metastable P alloys require careful thermomechanical processing to achieve

410

8 Titanium, Zirconium, and Hafnium

the required final microstructure. Controlled processing involves, first the worked or recrystallized condition, and then, if recrystallized, the grain size. Under certain melting conditions, the structure which occurs in an ingot of a metastable /? alloy ranges from small equiaxed grains at the surface, to elongated columnar grains, to large equiaxed grains at the center of the ingot. Recently, it was shown that there is a supra-transus "processing window" through which the alloy can be taken to result in a final fine equiaxed /? grain structure (Froes et al., 1985). This "processing window" is relatively wide for the lean metastable /? alloys and for large amounts of deformation. However, it is much more constricted for the richer alloys and for smaller amounts of deformation, making control of the /? grains much more difficult in the richer alloys. The mechanism by which the restoration to a low strain condition occurs is suggested schematically in Fig. 8-3 (Froes et al., 1985). In general, a fine /? grain structure is promoted by working below the transus temperature, followed by heating through the transus. Recrystallization follows the typical sigmoidal behavior, which is a function of temperature and prior deformation. The rate of grain boundary migration decreases inversely with annealing time, indicating a concurrent recovery process obeying second order kinetics. Grain growth follows the relationship

D1/n-Di'n =

(8-1)

where D is the grain size after annealing at a certain temperature for a time t, Do is the apparent initial grain size at t = 0, and n and A are constants. Generally, the kinetics of grain growth are not influenced by the prior grain size or amount of deformation, unless both recovered and recrystallized grains are measured. In this case, a

critical growth phenomenon occurs at low deformation levels (7-12%) (Froes et al., 1985). a Morphology The processing route determines the a morphology within the /? matrix, which can vary quite considerably; this morphology in turn strongly influences the mechanical properties (Froes et al., 1985). Two basic processing options are available: (1) /? processing, which is carried out either entirely above the /? transus or is completed below the /? transus but at a temperature high enough that very little a phase is present, and (2) a-/J processing, carried out below the /? transus temperature in the presence of a phase. Subsequent annealing below the p transus within about 175°C of the transus temperature results in a distribution of primary a phase which is related to the processing sequence and annealing temperature. With material generated by /? processes, a lenticular a morphology is achieved and maintained, while with a-j8 processing, the primary a P becomes globular during subsequent heat treatment (Fig. 8-4) (Kear, 1986; Joseph and Froes 1988). The change in morphology of a from lenticular to globular is a direct result of the prior deformation of the a phase. The strain energy in the a phase causes it to recrystallize and relax to a lower-surfaceenergy globular configuration. The transformation of lenticular a to globular a is a function of the annealing temperature and time and the amount of working the a phase has received, i.e., lightly worked a phase remains more lenticular than heavily worked a phase. Strength is virtually unaffected by the shape of the primary a but other properties such as fracture toughness and elevated-

STARTING STRUCTURE

DEFORMATION

RESTORATION INITIATED

RESTORATION PROGRESSES

AFTER ANNEALING

LOW TEMPERATURE DEFORMATION MIXED GRAIN STRUCTURE

STRAIN TENDS TO 8E LOCALIZED IN SANDS, CONCENTRATING IN SMALL GRAINS

RECRYSTALLIZATlOfTbc CUR S IN HEAVILY DEFORMED REGIONS

NEW GRAINS GROW INTO ADJACENT STRAINED REGIONS

PRIOR SMALL GRAINED REGIONS RETAIN SMALL GRAIN SIZE. MIXED STRUCTURE RESULTS*

DYNAMIC RECOVERY

RE CRYSTALLIZATION OCCURS

NEW GRAINS GROW INTO

UNIFORM FINE GRAIN

GIVES SMALL SU8-GRAIN

UNIFORMLY THROUGHOUT

ADJACENT STRAINED

STRUCTURE RESULTS

STRUCTURE AND UNIFORM

MATERIAL

REGIONS

RECRYSTALLIZATION OCCURS IN HEAVILY DEFORMED REGIONS, ESPECIALLY AT GRAIN BOUNDARIES

NEW GRAINS GROW INTO ADJACENT STRAINED REGIONS

INTERMEDIATE ("WINDOW") TEMPERATURE DEFORMATION MIXED GRAIN STRUCTURE

OEFORMATION, MINIMAL GRAIN GROWTH

HIGH TEMPERATURE DEFORMATION

MIXED GRAIN STRUCTURE

*

DYNAMIC RECOVERY GIVES LARGE SUB-GRAIN STRUCTURE AND UNIFORM DEFORMiATlON. GRAIN GROWTH OCCURS

RECRYSTALLIZATION MAY BE OYNAMIC AT T H E LOW TEMPERATURE BECAUSE OF THE EXTREMELY HIGH LOCALIZED STRAINS, IS EITHER METADYNAMIC OR STATIC AT INTERMEDIATE

t

GRAJN GROWTH RETARDED BY LOW STRAIN OF RECOVERED REGIONS. MIXED GRAIN STRUCTURE RESULTS 1

IN THE HIGH TEMPERATURE CASE OF THE ORIGINAL FINE GRAINS.

AND HIGH

TEMPERATURES,

THE LOCATION OF T H £ FINE GRAINS DOES NOT NECESSARILY

COINCIOE WITH THE LOCATION

Figure 8-3. Restoration process by which the strain in the titanium grains is reduced during an annealing treatment (Froes etal., 1985).

po k)

412

8 Titanium, Zirconium, and Hafnium

l

ular), see Sec. 8.2.4.1. A similar trend occurs with fatigue crack growth rate. However, optimum superplastic forming and diffusion bonding is found in material with a globular microstructures while creep performance is favored by lenticular a phase. Low cycle fatigue behavior is optimized with a globular a morphology. In P alloys, thermomechanical processing affects not only the microstructure, but also the decomposition kinetics of the metastable /? phase during aging. The increased dislocation density after working /? alloys leads to extensive heterogeneous nucleation of the equilibrium /? phase, which can suppress the formation of the brittle co (omega) phase (Froes et al., 1985).

8,2.3 Processing/Fabrication 8.2.3.1 Extraction

Figure 8-4. Microstructure of T i - 6 A l - 2 S n - 4 Z r 2Mo: (a) ft worked followed by an a-/? anneal to produce a lenticular a morphology; (b) a-/? worked and a-/? annealed to give predominantly an equiaxed a shape; and (c) OL-/3 worked followed by a duplex anneal: just below the p transus temperature (reduced volume fraction of equiaxed a compared to (b), and significantly below the /? transus temperature (to form the lenticular a between the equiaxed regions) (Kear, 1986; Joseph and Froes, 1988).

temperature flow characteristics (particularly as applied to creep, superplastic forming and diffusion bonding) are strongly influenced. High fracture toughness is associated with the a phase having a high aspect ratio (i.e., lenticular), while lower fracture toughness values at the same strength level correspond to the a phase having a low aspect ratio (i.e., glob-

The commercial production of titanium metal is based on the chlorination of rutile (TiO2) in the presence of coke or other form of carbon. The most important chemical reaction involved is TiO2(s) + 2Cl 2 (g) + 2C(s) -> (1) -> TiCl4(g) + 2CO(g) The resulting TiCl4 ("tickle") is purified by distillation and chemical treatments and subsequently reduced to titanium sponge using either Mg (Kroll process) or Na (Hunter process). The basic reaction of the Kroll process is (2) 2Mg(l) + TiCl4(g) Ti(s) + 2MgCl 2 (l) With either process the sponge produced is vacuum-distilled, swept with an inert gas, or acid leached to reduce the remnant salt content. A number of alternative processes have been evaluated for sponge production, including electrolytic,

8.2 Titanium

molten salt and plasma processes but none have reached full commercial status (Froes et al., 1985). 8.2.3.2 Ingot The titanium for ingot production may be either titanium sponge or reclaimed scrap. In either case, stringent specifications must be met for control of ingot composition. Modern melting techniques remove volatile substances from the sponge, so that ingot of high quality can be produced regardless of the method used for sponge production. However, for critical aerospace use, especially in engines, melting must be carried out to virtually eliminate the various types of defects. This has led to the development of melting techniques in which the time-temperature range where the metal is molten is increased (e.g., by electron beam and plasma techniques) compared to that of conventional vacuum arc consumable-electrode methods (Froes etal., 1985). Recycling of titanium scrap (revert) is an important aspect of cost effective production of titanium products. The revert which is recycled includes cut sheet, reject castings, machine turnings and chips. In 1988 almost 18 million kg of revert was used by U.S. producers. The normal melting practice for ingot production is double melting in an electricarc furnace under vacuum, which is considered necessary for all applications to ensure an acceptable degree of homogeneity in the resulting product. Triple melting is used to achieve better uniformity, and to reduce oxygen-rich or nitrogen-rich inclusions in the microstructure to low levels. Processes other than consumable-electrode arc melting are used in some instances for first-stage melting of ingot for noncritical applications.

413

8.2.3.3 Castings Castings are an attractive approach to the fabrication of titanium components since this technique allows production of relatively low cost parts (Froes et al., 1985). Basically a near net shape is produced by allowing molten titanium to solidify in a graphite or ceramic mold. Use of a ceramic mold, generally produced by the "lostwax" process, allows production of large, relatively high integrity, complex shapes. Enhanced mechanical properties in combination with increased size and shape-making capabilities, have resulted in greatly increased use of titanium castings in both engine and airframe applications. The shipment of titanium castings has increased by a factor of three over the past 15 years to a level of about 400000 kg/y. 8.2.3.4 Powder Metallurgy A number of powder metallurgy (PM) approaches have been evaluated for the titanium system including the blended elemental (BE), pre-alloyed (PA), rapid solidification, mechanical alloying (MA) and vapor deposition (VD) techniques (Froes et al., 1985; Froes and Eylon, 1990a; Suryanarayana et al., 1991; Froes and Suryanarayana, 1993). Using a press-and-sinter technique, the BE approach allows fabrication of low cost components from elemental and/or master alloy additions. However, because of the porosity resulting from this method, a result of the inherent salt, generally initiation-related properties such as S/N (stress/number of cycles to failure) fatigue are inferior to cast and wrought products. The PA approach yields mechanical properties at least equivalent to those of ingot products. However, less than desirable cost advantages, in combination with

414

8 Titanium, Zirconium, and Hafnium

a fear of the PM approach by design engineers have resulted in few applications. The powder metallurgy/rapid solidification (PM/RS) technique permits extension of alloying levels and much more refined microstructures than are possible using the ingot metallurgy (IM) technique. The greatly increased chemistry/microstructure "window" can lead to enhanced mechanical and physical properties in a variety of metallic systems. An alternative to PM/RS is mechanical alloying (MA) in which heavy working of powder particles results in intimate alloying by repeated welding and fracturing. This technique allows dispersoids to be produced, solubility extension, novel phase production and microstructural refinement. Production of alloys directly from the vapor grants even greater flexibility in microstructural development than RS or MA (Froes et al., 1995; Ward-Close and Froes, 1994). A semicommercial scale electronbeam vapor-deposition process has been constructed to produce alloys which are not obtainable by ingot methods or even rapid solidification. One example is the production of low density Ti-Mg alloys. Magnesium boils below the melting point of titanium, making fabrication of a liquid alloy impossible by conventional methods. 8.2.3.5 Joining Adhesive bonding, brazing, mechanical fastening and diffusion bonding are all used routinely and successfully to join titanium and its alloys (Froes et al., 1985). Welding methods of various types, including tungsten inert gas (TIG), electron beam and plasma, are also used very successfully with titanium and its alloys. In all types of welds, contamination by interstitial impurities such as oxygen and nitrogen

must be minimized to maintain useful ductility in the weldment. Thus, welding must be done under strict environmental controls to avoid pickup of interstitials that can embrittle the weld metal. 8.2.3.6 Wrought Product Processing This section addresses the primary processing of wrought (ingot) products to mill products. The following section will then be concerned with forming of these mill products to final components. Mill products include billet, bar, plate, sheet, strip, foil, extrusions, tubing and wire. Besides the reduction of section size and shaping, another objective of primary processing is the control (generally refinement) of the microstructure to optimize final mechanical property combinations (Froes et al., 1985; ASM Hdb., 1990). In general terms, as processing proceeds the temperature of the processing is decreased, with the /? transus temperature - below which the a phase can be present - being the critical temperature for the control of the microstructure. In many cases titanium is processed on the same equipment used for steel, with appropriate special auxiliary equipment. Other concerns in connection with the processing of titanium alloys include the high reactivity of titanium at elevated temperatures and the strain rate sensitivity; especially for the {$ alloys strength decreases as the strain rate is reduced. Billet production from an ingot starts above the /? transus temperature and proceeds at progressively decreasing temperatures. In some cases the /J grain size is reduced by a recrystallization treatment well above the /? transus temperature, however, the elimination of grain boundary a and the refinement of transgranular a necessitate working below the /? transus temperature.

8.2 Titanium

Bar, plate, sheet and foil products are produced on a relatively routine basis. In general, the processing is done at high temperatures, although the very high ductility of the metastable j8 alloys allows finishing of strip and foil by cold rolling. Forging is a very common method for manufacturing titanium alloy components. It allows both control of shape and manipulation of microstructure and hence of mechanical properties. Generally titanium alloys are considerably more difficult to forge than aluminum alloys and alloy steels, particularly when processing at temperatures below the /? transus is desirable. Extrusion, tubing and wire titanium products are also fabricated routinely, with the same caveats regarding microstructural control as for the product forms discussed above. 8.2.3.7 Wrought Product Forming For final use, mill products are formed to desired configurations with the same concerns regarding microstructural control as discussed in the previous section. Examples of forming of wrought product include isothermal/hot forging, sheet metal forming, and superplastic forming/ diffusion bonding. Isothermal and hot forging are special forging operations in which the die temperatures are close to the metal temperature, i.e., much higher than in conventional forging. This reduces chill effects and allows close to net shape production. Strain rates are much lower than normal, contributing to the near net shape capability. The metastable /? alloys, with a low P transus temperature, are particularly amenable to the isothermal forging process. Sheet metal forming is conducted either under hot conditions, which generally al-

415

low larger, more precise amounts of deformation, or under cold conditions, giving rise to lower costs. Hot forming of titanium alloys is conducted in the range 595815°C with enhanced formability and reduced spring back. Formability increases with increasing temperatures but at higher temperatures contamination can become a problem, sometimes necessitating an inert atmosphere or a coating. /? alloys are easier to cold form than a and a-/? alloys. The high degree of spring back exhibited by titanium alloys sometimes requires hot sizing after cold forming. This reduces internal stresses and restores compressive yield strength. Superplastic forming/diffusion bonding makes use of the fact that fine grained material can deform extremely large amounts, especially at very low strain rates (0.0001 to 0.01s" 1 ). Superplastic forming (SPF) is the propensity of sheet material to sustain very large amounts of deformation without unstable deformation (tensile necking); for example, fine grained (<10 |im) Ti-6 A1-4V can be deformed > 1000% in tension at 927 °C. Diffusion bonding (DB) is a solid state bonding process in which a combination of pressure and temperature allows production of a metallurgically sound bond. Superplastic forming, commonly under gas pressure, is now used routinely as a commercial sheet metal fabrication process for reduced cost, and complex shapes. The combined SPF/DB process has seen little use, predominantly because of problems in inspecting the integrity of the bond region. 8.2.3.8 Machining Previous sections of this chapter discussed a number of approaches to reduce the cost of titanium components, particularly near-net shape methods. However

416

8 Titanium, Zirconium, and Hafnium

most titanium parts are still produced by conventional means involving substantial machining (Froes et al., 1985). As a result, machining of titanium and its alloys has been extensively evaluated and well-defined procedures for various types of machining operations have been specified, including turning, end milling, drilling, reaming, tapping, sawing, and grinding. In many instances considerable amounts of machining are required for the production of complex components from mill products such as forgings, plate and bar, i.e., a high buy-to-fly ratio. Titanium is chemically reactive, leading to a tendency to weld to the tool and chipping and premature failure. Other problems involve the low heat conductivity of titanium, which adversely affects tool life, and the ease of damaging the titanium surface. The latter effect is of particular concern because surface integrity strongly influences crack initiation related properties such as fatigue. The machining of unalloyed titanium is similar to \~\ hard austenitic stainless steel. High quality sharp tools, carbides for high productivity and high speed tool steels for more difficult operations are required for titanium. This, in combination with slow speeds, heavy feeds and the correct cutting fluids, generally results in good machining behavior for titanium. Cutting fluids recommended are oil-water emulsions and water soluble waxes at high cutting speeds, and low viscosity sulfurized oils and chlorinated oils at low speeds; in all cases the cutting fluids should be removed after machining, especially before heat treatment, to avoid potential stresscorrosion cracking problems resulting from the use of chlorinated oils.

develop engineered metals including metal matrix composites (MMCs). The so-called CermeTi family of titanium alloy matrix composites, fabricated using the blended elemental approach, incorporate particulate ceramic (TiC or TiB2) or intermetallic (TiAl) as a reinforcement (Froes and Suryanarayana 1993). These composites exhibit minimal particle/ matrix interaction while maintaining the integrity of the essentially 100% dense homogeneously dispersed particles within the matrix. Examples of CermeTi material are shown in Fig. 8-5. An innovative method (XD) for the production of in situ discontinuous titaniumbased composites has resulted in interesting property combinations in alloys such as the intermetallic y (Christodoulou and Brupbaker, 1990), but to date there are no applications. Reinforcement with continuous ceramic composites enhances the strength and modulus of terminal titanium alloys, such as Ti-6A1-4V (MacKay et al., 1992; Upadhyaya et al., 1994). However, control of the reaction zone between the fiber and the matrix, inferior transverse properties, and cost remain major concerns. Innova-

8.2.3.9 Metal Matrix Composites The success exhibited by organic matrix composites has led to parallel efforts to

Figure 8-5. Microstructure of CermeTi material, TiC reinforcing particles (courtesy of Dynamet Technology, Inc.).

417

8.2 Titanium

tive fabrication techniques such as plasma spray deposition and electron beam vapor deposition may help in controlling cost, while the design lessons learned with nonisotropic polymeric composites should be applicable in engineering metal composite structures. The potential weight savings obtainable by replacing much heavier superalloys in both engine and airframes has resulted in considerable work being conducted on a2 and y MMCs (Froes, 1991). Recently, increased attention has been given to the richer Nb varieties of Ti 3 Al-Nb known as the "orthorhombic" alloys (22-27 at.% Nb) as matrix materials (Froes et al., 1992; Froes, 1994). 8.2.4 Mechanical Properties 8.2.4.1 Cast and Wrought Terminal Alloys

The mechanical properties of titanium alloys not only depend on the chemistry but are also strongly influenced by their microstructure, the latter in turn being dependent on the processing conditions. The tensile properties of selected cast and wrought terminal titanium alloys are summarized in Table 8-3 (Polmear, 1989). The influence of a morphology was discussed in Sec. 8.2.2.5, where it was pointed out that a lenticular shape favors high fracture toughness (Klc) while a globular morphology optimizes ductility. The effect of a morphology and section size on tensile properties and fracture toughness are demonstrated in Tables 8-4 and 8-5 (ASM, 1990) and illustrated in Fig. 8-6 (Froes et al., 1985); as strength increases, fracture toughness decreases, and vice versa. Chemistry (in particular the interstitial content, e.g., O2) influences fracture toughness, with high values of Klc associated with low O 2 values; texture can also have an effect.

UTS (ksi)

120

140

160

180

200

160 140

800 900 1000 1100 1200 1300 1400 1600 UTS (MPa)

Figure 8-6. Variation of fracture toughness with ultimate tensile strength (UTS) for Corona5 (Ti-4.5 Al5Mo-1.5Cr). At a given strength level, a lenticular a gives a higher fracture toughness than a globular morphology (Froes et al., 1985).

The fatigue behavior of titanium alloys can be divided into S/N fatigue and fatigue crack growth rate (FCGR or da/dN versus AK, where a is crack length and N the number of cycles). Within S/N fatigue a further subdivision can be made between low cycle fatigue (LCF) and high cycle fatigue (HCF). For LCF, failure occurs in 104 cycles or less, while in HCF failure occurs at greater than 104 cycles. Different uses favor different techniques for determining LCF, specifically strain controlled and load-controlled tests, Table 8-6 and Fig. 8-7 (Donachie, 1988). Both notch concentration (Xt) and overall surface condition can strongly influence LCF. The beneficial effect of relatively gentle surface conditioning is shown in Fig. 8-8 (Donachie, 1988); more severe working of the surface can result in the formation of

Table 8-3. Compositions, relative densities and typical room temperature tensile properties of selected wrought titanium alloys. Common designations

a alloys CPTi99.5%IMI115, Ti-35A CPTi99.0%IMI155, Ti-75A IMI260 IMI317 IMI230 Near-& alloys 8-1-1 IMI679 IMI685 6-2-4-2S Ti-11 IMI829

Al

5

Sn

Zr

Mo

V

Si

Other

Relative density

Conditionsa'b

0.2% Proof stress (MPa)

0 0 0.2 Pd

4.51 4.51 4.51 4.46 4.56

annealed 675 °C annealed 675 °C annealed 675 °C annealed 900 °C ST (a) duplex aged 400 and 475 °C

170 480 315 800 630

240 550 425 860 790

25 15 25 15 24

4.37 4.82 4.49 4.54 4.45 4.61

Annealed 780 °C ST(a + £) aged 500 °C ST($ aged 550 °C ST(a + jff) aged 590 °C ST(jff) aged 700 °C ST08) aged 625 °C

980 990 900 960 850 860

1060 1100 1020 1030 940 960

15 15 12 15 15 15

4.46

annealed 700 °C ST(a + P) aged 500 °C ST(a + jS) aged 500 °C ST(oc + £) aged 500 °C ST(a + £) aged 550 °C ST(0) annealed 590 °C ST(a + £) aged 500 °C annealed 700 °C

925 1100 1000 1190 1170 1170 1200 860

990 1170 1100 1310 1275 1270 1310 945

14 10 14 15 10 10 13 15

ST(0) aged 480 °C ST(0) duplex aged 480 and 600 °C ST(0) aged 580 °C ST(0) aged 540 °C ST(0) aged 580 °C

1200 1315 1240 1280 1130 1250

1280 1390 1310 1400 1225 1320

8 10 8 6 10 8

2.5 2.5 C u

Tensile Elongastrength tion (MPa) (%)

IMI550 IMI680 6-6-2 6-2-4-6 IMI551 Ti-8Mn p alloys 13-11-3 8-8-2-3 Transage 129

PC 10-2-3 a

CD

E" 3 N o c" 13 Z3

8 2.25 6 6 6 5.5

11 2 2 3.5

5 5 4 1.5 3

1 1 0.5 2 1 0.3

1 0.25 0.25 0.2 0.1 0.3

0.35 Bi 1 Nb

(x-fi alloys

IMI318, 6-4

oo

6 4 2.25 6 6 4

4 2 11 2 2 4

4

0.5 0.2

4 4 6 6 4

0.7(Fe,Cu)

0.5 8Mn

3 3 2 3 3

4.5

6

2

11 4

11.5 8 4

13

11 Cr

8 11 8 10

2Fe 6Cr 2Fe

4.60 4.86 4.54 4.68 4.62 4.72 4.87 5.07 1 4.85 j 4.81 4.82 4.65

ST (a), ST (a + /?), ST (/?) correspond to solution treatment in the a, a + /?, and /?-phase fields, respectively; b annealing treatments normally involve shorter times than aging treatments.

Q.

IE

i

419

8.2 Titanium

Table 8-4. Yield strength and plane strain fracture toughness of various titanium alloys. Alloy

a Morphology or processing method

Ti-6A1=4V

Yield strength

equiaxed transformed oc-p rolled + mill annealeda

(MPa)

Plane-strain fracture toughness (XIC) (MPa ^ m )

910 875 1095

44-66 88-110 32

Ti-6Al-6V-2Sn

equiaxed transformed

1085 980

33-55 55-77

Ti-6Al-2Sn-4Zr-6Mo

equiaxed transformed

1155 1120

22-23 33-55

OL + P forged, solution treated and aged p forged, solution treated and aged

903

81

895

84

oc-p processed P processed

1035-1170 1035-1170

33-50 53-88

Ti-6Al-2Sn-4Zr-2Mo forging

Ti-17 Standard oxygen (~0.20 wt.%)

Table 8-5. Relation of tensile strength of solution-treated and aged titanium alloys to size. Alloy

Tensile strength (MPa) of square bar in the section size of: 13 mm

25 mm

50 mm

75 mm

100 mm

150 mm

1105 1205 1170 1170 1240 1310 1310 1310

1070 1205 1170 1170 1240 1310 1310 1310

1000 1070 1170 1170 1240 1310 1310 1240

930 1035 1140 1105 1240 1310 1310 1240

1105 1105 1170 1310 1310 1170

_ 1105 1170 1310 1170

Ti-6A1-4V Ti-6Al-6V-2Sn(Cu + Fe) Ti-6Al-2Sn-4Zr-6Mo Ti-5Al-2Sn-2Zr-4Mo-4Cr (Ti-17) Ti-10V-2Fe-3Al Ti-13V-llCr-3Al Ti-ll.5Mo-6Zr-4.5Sn (P III) T i - 3 A I - 8 V - 6 C r - 4 Z r - 4 M o (p C)

Table 8-6. Strain control low-cycle fatigue life of Ti-6242Sat480°C. Test fre-

Total

Number of cycles to failure

(cycles/ min)

range (%)

Acicular structure

Equiaxed a structure

0.4 10 0.4 10

1.2 1.2 2.5 2.5

1 196 3 715 273 353

10 5003 31 000a 722 1 166

a

Run out.

cracks and degraded LCF behavior. The effect of Kt and crack propagation on LCF life on preloaded Ti-6A1-4V at 205 °C is shown in Fig. 8-9 (Donachie, 1988). Surface condition can also strongly influence HCF (Fig. 8-10) (Donachie, 1988). The fatigue endurance limit is relatively flat to at least 315°C, Fig. 8-11 (Donachie, 1988), with benefits apparent for titanium alloys over steels. The FCGR performance generally parallels fracture toughness, with the reserva-

420

8 Titanium, Zirconium, and Hafnium

16

160

-

12

I 120

4

40

Peened or Machined

©

U

I 103

I I I I 104

I

I I I I I 1 1 1 ! 10s 10* Cycles to Failure

I

I

I I I 107

Figure 8-10. Effect of surface finish on LCF (Donachie, 1988). Beta Forged

10% Primary Alpha

50% Primary Alpha

Figure 8-7. Low cycle fatigue (LCF) life of Ti-6 Al4V with different structures (Donachie, 1988).

0

100

Temperature, °C 200 300 400

500

600

R=-l K,= l Lathe 1\irned

2 400 Ti-6A1-4V Ti-8Al-lMo-lV

Shot Peened

Glass Bead Blasted

I 200

J 40

I

300

100

J

1

500

Cyclic Life-95

Figure 8-8. Effects of surface condition on LCF life of Ti-6 Al-4 V at 21 °C (Donachie, 1988).

60

First Crack

40

U 20

103

104

10s

Cycles

Figure 8-9. LCF life strength of Ti-6 Al-4 V at 204 °C, showing the effect of Kt (notch concentration) and crack propagation rate on life (Donachie, 1988).

200

400

600 8 Temperature~°F

Figure 8-11. High cycle fatigue (HCF) at 5 x 107 cycles, showing the benefits of titanium over steel (103 in ~ 25 m). R: ratio of the minimum to maximum stress (Donachie, 1988).

tion that severe corrosive environments (such as 3.5% NaCl solution) can adversely affect the FCGR by an order of magnitude. An example of the strong influence of the microstructure on FCGR for the Ti-6 Al-4 V alloy is shown in Fig. 8-12 (Donachie, 1988); the /^-annealed condition proved to be significantly better than the mill annealed material. The a and near-a alloys generally exhibited superior high temperature behavior (Fig. 8-13) (ASM, 1990). The reason why

8.2 Titanium Stress-intensity factor range, AK, ksi • in. 1 / 2 10

20

40

60 80 100 10~ 2 B C

10" 1

10" 10" 2

.2 |

.£

1

10~4

•a

10" 3

10" b

Room temp air R = 0.3 1 Hz L-T orientation 15.9-mm (0.62-in.) plate 0.2% wt% oxygen

10-

10"

I

10

i

20

i

40

£

these alloys have replaced steels in advanced jet engines is clearly demonstrated in Fig. 8-14 (ASM, 1990); titanium alloys are now used up to about 600 °C. Generally these alloys contain Si for enhanced creep behavior, Figs. 8-15 and 8-16 (ASM, 1990). Titanium alloys have good cryogenic properties, a alloy Ti-5 Al-2.5Sn and the a-/? alloy Ti-6A1-4V finding extensive use. A number of /? alloys have been developed over the years, with current activity being concentrated on the alloys shown in Table 8-7. In general, these alloys exhibit higher strength-toughness combinations than the a-ft alloys such as Ti-6A1-4V (Boyer and Hall 1993).

8.2.4.2 Iiitermetallic Alloys

A - Mill annealed B - Diffusion bond thermal cycle C - Beta annealed 10"

10-6

Intermetallics are discussed in detail by Sauthoff in Chapter 11 of this book. Thus only a brief description will be given here, with emphasis on indicating how both chemistry and processing/microstructure can be adjusted to control mechanical properties.

10i

i

60

80 100

Stress-intensity factor range, AK, MPa • m 1 / 2

Figure 8-12. Effect of microstructure on the fatigue crack growth rate (FCGR) for the Ti-6 Al-4 V alloy (Donachie, 1988).

Ti834 TA5E

421

Ti-6AI-2Sn-4Zr2Mo-0 8Si

Ti-6AI-2Sn-4Zr-2Mo

IMI 6 8 5

7117 Betacez

Heat treatment capacity

a alloys

Near-a alloys

a + p alloys

Near-p alloys

p alloys

Figure 8-13. Major characteristics of the three terminal classes of titanium alloys (ASM, 1990).

422

8 Titanium, Zirconium, and Hafnium Temperature, °F 390

212

570

750

1110

930

250

1000

200

I •s 100

300 Temperature, °C

600

Figure 8-14. Specific tensile strength (strength divided by density) of various titanium alloys, compared with steels once used in gas turbine engines (ASM, 1990).

Table 8-7. Typical /? titanium alloy compositions.

Alloy

10-2-3 15-3 I321S £111 a

Composition Al

Sn

Zr

V

Fe

Mo

Cr

2 3 3 3 -

— 3 4 4.5

— -

10 15 8 -

3 -

— 4 15 11.5

— 3 6 -

-

•

6

Nb

2.6

Si

0.2

Ti bal.1 bal. bal. bal. bal.

bal. = balance.

Typical properties of the Ti 3 Al(a 2 ) type of titanium aluminide are shown in Table 8-8 (Froes et al., 1991, 1992; Froes, 1994). A good example of how microstructural control can lead to tailoring of the mechanical properties is the very high ambient temperature ductility obtained by spe-

cial processing to produce an optimum amount of equiaxed primary a 2 . Recently, control of the microstructure in TiAl(y) alloys, especially those in the two-phase y + a2 phase field, has led to interesting mechanical property combinations including an ambient temperature

423

8.2 Titanium Larson-Miller time (/in hours) - temoerature (°F) parameter, P 30

31

32

33

34

35

500

P = (Tin °F + 460) (20 + log t) x 10 P= (Tin °C + 255.6) (20 + log /) x 10

17

18

19

Larson-Miller time (/in hours) - temperature (°C) parameter, P

Figure 8-15. Creep behavior of three elevated temperature titanium alloys; all three alloys contain Si for enhanced performance. The alloys were either a-/? or ^-processed (ASM, 1990). Table 8-8. Typical properties of Ti3Al-type titanium aluminidesa. Alloy

UTS (MPa)

YS (MPa)

Elong. (%)

Klc (M°Pa y/m)

Creep rupture b

Ti-25A1 Ti-24Al-llNb Ti-25Al-10Nb-3V-lMo Ti-24Al-14Nb-3V-0.5Mo Ti-24.5Al-17Nb

538 824 1042 1010 940 1133 963

538 787 825 952 705 989 860

0.3 0.7 2.2 26.0c 5.8 10.0 3.4 6.7

— 13.5

_ 44.7 360 62 476 0.9

Ti-25Al-17Nb-lMo Ti-15Al-22.5Nb Compositions in at.%;

b

hours at 65O°C/38 MPa;

ductility of almost 5% in a Ti-46.5A12.5V-lCr alloy. Lenticular microstructures result in high toughness/low ductility, while the duplex microstructures give the opposite combination of these properties (Fig. 8-17 and Table 8-9) (Froes etal., 1991, 1992).

c

28.3 20.9 42.3

specially processed.

8.2.4.3 Cast Alloys Because cast alloys are produced in "near-net" shape directly from the molten state they inherit a microstructure which cannot be modified by the thermomechanical processing used with cast and wrought

424

8 Titanium, Zirconium, and Hafnium Approximate temperature, °F

500 700

600

700

800

900

1000

110 —

600

\Ti- 6AI-2Sn-4Z ^-6Mo 500

- 80 w -

\ V

Ti-8AI-1Mo-1V

\

"—6 0

400

\

300

\Ti-6AI-2Sr v4Zr-2Mo

5AI-2.5Sn

40

o

\

200

\

\

Ti-6AI-6V-2Sn

\

100

>^

\

-

20

\

Ti-6AI-4V

n 260

n

315

a)

370

425

480

535

590

Approximate temperature, °C

Temperature. °F 930

840

1110

1020

87

600

72.5

500

^ IMI 829 \

400

\ \

300

\

^ \

43.5

\ \ ^

6 200 o

58

IMI 834

\ IMI 685

\

29

\ Ti 6242S

100 450

b)

500 550 Temperature, °C

(ingot) material (Froes et al., 1985). Additionally a number of defects can occur in castings, such as porosity, which can impair mechanical properties. The microstructure of cast products, for example in the Ti-6 A l - 4 V alloy, consists of large /? grains, extensive grain boundary a, and elongated coarse intragranular a

600

14.5

Figure 8-16. Comparison of the creep strengths of various titanium alloys (ASM, 1990).

which can occur in colonies of similarly aligned plates or in a Widmanstatten morphology. Thus characteristics such as strength, fracture toughness, fatigue crack growth rate and creep behavior are at relatively high levels (Table 8-10) (Donachie, 1988). However ductility and S/N fatigue are lower than for cast and wrought prod-

8.2 Titanium

425

Figure 8-17. Variety of microstructures in Ti-47.5Al-2.5V-lCr (y), see Table 8-9 (Froes et al, 1991, 1992)

Table 8-9. Microstructure and mechanical properties in Ti-46.5A1^2.5V-lCr TiAl-type titanium aluminide a.

YS (MPa) UTS, (MPa) Elong. (%) K I c (MP a v /m)

FL

NL

DM

NG

360 400 0.5 21

430 480 2.3 17

440-450 505-538 3.3-4.8 12

387 468 1.7 12

a

compositions in at.%. FL, fully lamellar; NL, near lamellar; DM, duplex; NG, near gamma.

ucts (Fig. 8.18) (Froes et al., 1985). Both ductility and S/N fatigue can be enhanced by use of either innovative heat-treatments

or the use of hydrogen as a temporary alloying element (thermochemical processing, TCP) to refine the microstructure (Fig. 8-19) (Froes et al., 1990b). The high cycle fatigue of alloys such as Ti-6 A1-4V can be enhanced by hot isostatic pressing (HIP'ing) (Froes and Hebeisen, 1994). Casting of titanium alloys other than the conventional Ti-6A1-4V alloy is also possible. An example is the Ti-3 Al-8 V 6 C r - 4 Z r - 4 M o alloy (38-6-44 or Beta C), which exhibits excellent tensile properties and impressive fatigue behavior, with an endurance limit 85 % above the average value typical of the Ti-6A1-4V alloy (Froes et al., 1985).

426

8 Titanium, Zirconium, and Hafnium

Table 8-10. Typical room-temperature tensile properties of several cast titanium alloys. Conditions

Tensile strength (MPa)

Yield strength (MPa)

Elongation

Reduction in area

as-cast or annealed as-cast or annealed duplex annealed annealed

550 1035 1035 805

450 890 895 745

17 10 8 11

32 19 16 -

Alloy

Commercially pure titanium Ti-6A1-4V Ti^6Al-2Sn-4Zr-2Mo Ti-5Al-2.5Sn-ELI

~

!

'I

! —

TI-6AI-4V SMOOTH AXIAL FATI6UE R= - 0 . 1 ROOM TEMP.

-

-160

-120

i

IUM STRESS,

1200

1

-

J

400

- 40

1

L_

1

103

104

105

1

10B

Figure 8-18. Room temperature smooth axial fatigue versus maximum cyclic stress for Ti-6A1-4V. Data scatterbands for cast, cast and hot isostatically pressed, and annealed wrought (ingot) material (Froes et al., 1985).

CYCLES TO FAILURE. N |

i^::^xi^juJM f^^y-^ ---: -u. '4 r- f,~. '

Figure 8-19. Refinement of the microstructure of cast Ti-6 Al-4 V (left) by use of the thermochemical processing technique (right) (Froes and Eylon, 1990 b; Froes and Suryanarayana, 1993).

8.2.4.4 Powder Metallurgy Alloys The tensile properties of blended elemental (BE) products (elemental additions) meet typical minimum wrought specification properties (Table 8-11). However, because of the remnant salt (from the

extracation process, see 8.2.3.1) and associated porosity, fatigue behavior is below wrought levels. However, this behavior may be enhanced in a similar fashion to cast products by use of innovative heattreatments or TCP. At a cost penalty, properties may also be improved by using

427

8.2 Titanium

Table 8-11. Typical tensile properties of blended elemental Ti-6A1-4V compacts compared to mill-annealed wrought products. Material Cold isostatic press and HIP (CHIP) Press and sinter (no HIP'ing) Wrought mill annealing Typical minimum properties (MIL-T-9047)

0.2% YS (MPa)

UTS (MPa)

Elongation

Reduction in area

827 868 923 827

917 945 978 896

13 15 16 10

26 25 44 25

Table 8-12. properties of Ti-6A1-4V pre-alloyed powder compacts. 0.2% YS UTS (MPa)

(MPa)

Elongation (%)

930

992

15

Reduction in area (%)

Klc

33

77

% Strain

Table 8-13. Tensile data for Ti-1 Al-8V-5Fe high strength alloy. Alloy

Ti-lAl-8V-5Fe Conventional Ti-6A1-4V

TI6242S

Tensile properties YS (MPa)

UTS (MPa)

Elong.

1390

1480

895

965

8 14

higher priced salt-free titanium sponge or a newly available hydride powder produced by a calcium process (Moxson, 1993). The tensile and fracture toughness properties of pre-alloyed (PA) material are at levels at least equivalent to wrought products (Table 8-12) (Froes et al., 1990a; Froes and Suryanarayana, 1993), and, with adequate precautions to avoid contamination of the powder, S/N fatigue behavior is also at least at ingot levels. As with BE products, S/N fatigue can be further improved by use of innovative heat-treatments or TCP.

0.0

Figure 8-20. Creep properties (650 °C, 140 MPa) of rapidly solidified (RS) dispersion-strengthened alloys compared to a conventional wrought alloy. The base RS alloy is Ti-13.5Sn~3Al-lNb-2.5Zr-0.2Mo (Suryanarayana et al., 1991; Froes and Suryanarayana 1993).

Rapidly solidified (RS) titanium alloys containing rare earth additions show some improvement in creep behavior (Fig. 8-20) (Suryanarayana et al., 1991; Froes and Suryanarayana, 1993) but have not yet seen commercial use. The RS approach can also be applied to produce high strength alloys such as the normally segregation-prone Ti-1 A l - 8 V - 5 F e , Table 8-13 (Suryanarayana et al., 1992). Little advantage has been achieved for the intermetallics using RS, with the caveat that the near net shape processing may offer an advantage for the very difficult to fabricate y-compositions.

428

8 Titanium, Zirconium, and Hafnium

Development of mechanically alloyed (MA) titanium alloys is at a very early stage with virtually no mechanical properties available (Froes and Suryanarayana, 1993). Early indications, however, suggest that improved dispersions of second phase particles and enhanced strength-ductility combinations may occur, the latter in very fine grained nanostructured material. 8.2.4.5 Welded Components Welding generally increases strength and hardness, and decreases ductility (Froes et al., 1985). Welds in unalloyed titanium grades 1, 2, and 3 (Donachie, 1988) do not require postweld treatment unless the materials will be highly stressed in a strongly reducing atmosphere. Welding of the a class of alloys and leaner a-/3 alloys such as Ti-6 A1-4V can be accomplished with relative ease. However, welds in more /?-rich a-/? alloys such as Ti-6 A l 6 V-2Sn have a high likelihood of fracturing with little or no plastic straining. Weld ductility can be improved by postweld heat treatment consisting of slow cooling from a high annealing temperature. Rich /?-stabilized alloys can be welded, and such welds exhibit good ductility. However, here the aging kinetics of the weld metal may be substantially different from those of the parent metal. The intermetallic compounds Ti3Al and TiAl are difficult to weld because of their low inherent ductility and the microstructures which develop in the weld region. A range of energy inputs into the weld region using EB techniques allows control of heat input and elimination of solid state cracking in the fusion zone (Froes and Baeslack, 1993). 8.2.4.6 Machined Components The surface of titanium alloys may be damaged during machining and grinding;

and this damage can lead to a reduction in fatigue strength and stress corrosion resistance (Froes et al., 1985). Shallow compressive stresses can enhance fatigue behavior. 8.2.4.7 Metal Matrix Composites The various grades of CermeTi offer higher elevated temperature strength, increased hardness, and improved modulus over the monolithic titanium alloy while maintaining both the fracture toughness and machinability of a metal (albeit more difficult), as opposed to those of a brittle ceramic (Froes and Suryanarayana, 1993). The mechanical behavior of T i - 6 A 1 4 V reinforced with continuous SiC fibers is shown in Table 8-14 (Smith and Froes, 1984; Froes et al., 1985). Greatly enhanced specific strength is obtained in oc2-type titanium aluminide/SiC composites compared to conventional superalloys, with dramatic weight savings of up to 75 % by replacing a conventional disc and spacer assembly by a titanium aluminide reinforced ring (Driver, 1990). As a matrix, the richer orthorhombic a2 alloys exhibit increased ambient temperature ductility and enhanced oxidation resistance. However, they are more costly and of higher density (Froes et al., 1991, 1992; Froes, 1994). 8.2.5 Chemical Properties/Corrosion Behavior Titanium is used in aerospace and commercial applications because of its high strength-to-density ratio, good fracture characteristics, and general corrosion resistance. Titanium's excellent resistance to most environments is the result of its strong affinity for oxygen and tendency to form a stable, tightly adherent, protective surface film (Froes et al., 1985). This film consists basically of TiO 2 at the metal-en-

429

8.2 Titanium

Table 8-14. Mechanical properties of titanium metal matrix composites.

Table 8-15. Acid-concentration limits for ASTM gradesa 2, 7, and 12 titanium in pure reducing acids.

System

Elong. (GPa) La

Acid/

120 240

HC1 24 °C boiling

6 0.6

25 4.6

H 2 SO 4 24 °C boiling

5 0.5

48 7

10 1.5

H 3 PO 4 24 °C boiling

30 0.7

80 3.5

40 2

Ti-6A1-4V SCS-6(SiC)/Ti-6Al-4V L: Longitudinal;

b

UTS (MPa) La

UTS (MPa)

890 1455

890 340

Tb

T: transverse.

vironment interface with underlying thin layers of Ti 2 O 3 and TiO. It forms naturally and is maintained when the metal and its alloys are exposed to moisture or air. In general, anhydrous conditions such as provided by chlorine or methanol as well as uninhibited reducing conditions should be avoided. The passive film formed in air may not be adequately stable and may not be regenerated if it is damaged during exposure to these environments. General (uniform) corrosion rates for titanium and many of its alloys exposed to a wide variety of environments are available (Froes et al., 1985; Schutz, 1994). In general, commercial purity titanium is resistant to natural environments, including sea, fresh, brackish and mine waters, food products, crude oils, body fluids, and waste materials. The outstanding resistance of unalloyed titanium and the Ti~ 0.2Pd alloy (ASTM Grades 2 and 7 - see Table 8-15) in chloride-containing, aqueous environments is well-established. With few exceptions, unalloyed titanium performs well when exposed to oxidizing inorganic acids (e.g., nitric and chromic acid), aqueous ammonia, anhydrous ammonia, molten sulfur, pure hydrocarbons, aqua regia, hydrogen sulfide, wet chlorine, most organic acids, dilute caustic solutions, and chlorine dioxide. Titanium is not particularly resistant to pure reducing inorganic acids (i.e., those

Acid-concentration limitb (wt.%) Grade 2

Grade 7

Grade 12 9 1.3

a

Grade 2: Ti-50A; grade 7: Ti-0.2Pd; and grade 12: Ti-0.3Mo-0.8Ni; b for a corrosion rate of about 5 mil per year (about 0.13 mm per year).

that generate hydrogen during the metalacid reaction) such as sulfuric, hydrochloric, and phosphoric acids. The metal is dissolved rapidly by hydrofluoric acid. Other environments which should be avoided include fluoride-containing solutions (e.g., ammonium fluoride), hot concentrated caustics, certain organic acids, (e.g., oxalic, concentrated citric, trichloroacetic, and non-aerated boiling formic), and powerful oxidizing agents (e.g., anhydrous liquid and gaseous chlorine, liquid and gaseous oxygen, anhydrous red fuming nitric acid (RFNA), anhydrous nitrogen tetroxide, and liquid bromine). Powerful oxidizers are especially to be avoided because, under certain conditions such as impact, the reaction can be pyrophoric. Figure 8-21 (Froes et al., 1985) shows that the use of titanium can be extended into the "reducing acid" region by alloying the metal with small amounts of a noble metal such as 0.2 wt. % palladium, Table 8-15 (Froes et al., 1985). A similar but smaller benefit can be achieved by small alloying additions of nickel and molyb-

430

8 Titanium, Zirconium, and Hafnium

p

° i CHLORIDE

ACI

TANTALUM

RCONIUM LOYALLY B

i i inv

TITANIUM

' ,

t

HASTELLOY ALLOY C -H-MONEL HASTELLOY ALLOY F ZIRCONIUM 316 STAINLESS

N0.20AUflY

INCONEL ~« -*PX -

HASTELLOY MONEL

304 STAINLESS

— — OXIDIZING

REDUCING*—-

Figure 8-21. General corrosion behavior of commercially pure titanium and Ti-Pd alloys compared to other metals and alloys in oxidizing and reducing acids, with and without chloride ions. Each metal or alloy can generally be used for those environments below its respective solid lines (Froes et al., 1985).

denum (e.g., 0.3 wt.% Mo and 0.8 wt.% Ni; ASTM Grade 12 titanium). Unalloyed titanium is especially useful for applications where essentially no corrosion products can be tolerated in the process fluid. The metal is used extensively in the fabrication of food, drug, and dye processing equipment where even trace amounts of metal ion contamination could adversely affect the quality, color, and/or taste of the product produced. Titanium, in most environments, is an effective cathode, thus coupling titanium to a less noble metal can result in a high galvanic corrosion current and rapid dissolution of the anodic material, and the titanium may absorb hydrogen. Titanium and its alloys are susceptible to inhibitor type concentration-cell corrosion when, for example, oxidizing heavy metal ions are used to inhibit general corrosion and crevices exist. Concentrationcell corrosion of titanium can be mitigated in some cases by using either ASTM Grades 7 or 12 titanium (see Table 8-15) for fabrication of entire components or just for local crevice zones. Palladium and

nickel in these alloys provide improved passivity (i.e., anodic protection) in the crevices. Resistance to chloride-induced pitting attack is a primary reason for using titanium (e.g., replacing Type 316L stainless steel in petroleum refinery processes). However, under certain conditions, titanium is susceptible to pitting attack and has been reported to pit in the hot 130°C (270 °F) brine solutions in salt evaporators. As in the case with concentration-cell corrosion, pitting attack can also be mitigated by using ASTM Grades 7 and 12. Titanium alloys are susceptible to stresscorrosion cracking (SCC) in a number of environments, including anhydrous methanol containing trace quantities of halides, anhydrous RFNA, and hot chloride-containing salts. Several of the alloys and unalloyed titanium (containing relatively high oxygen contents) are known to fail in ambient temperature sea water if the materials contain pre-existing cracks. The strong influence of microstructure on SCC has been demonstrated in the metastable /? titanium alloy Beta III (Ti11.5Mo-6Zr-4.5Sn) (Guernsey et al., 1972). This work demonstrated that the alloy was susceptible to SCC when equiaxed /? grains and continuous grain boundary a were present, while a worked material in which no such microstructural features were observed was immune to SCC. Although there have been no known service failures related to hot salt stress-corrosion cracking (HSSCC), HSSCC presents a potential limitation to the long duration exposure of highly stressed titanium alloys at temperatures above about 220 °C. The near immunity of relatively high strength titanium alloys to corrosion fatigue in chloride-containing solutions allows these materials to be used in many

8.2 Titanium

hostile environments (e.g., body fluids) where other alloys have failed when subjected to cyclic stresses. The tenacious passive film which forms naturally on titanium and its alloys provides excellent resistance to erosion corrosion. For turbine-blade applications where the components are impinged by high-velocity water droplets, unalloyed titanium has been shown to have superior resistance compared to conventional blade alloys (e.g., austenitic stainless steels and Monel). It is known that the fatigue behavior of titanium and its alloys is surface condition sensitive; surface damage by fretting can adversely affect the ability of these materials to withstand cyclic stress. For example, fretting corrosion can reduce the fatigue strength of a titanium alloy, such as T i 6A1-4V, by more than 50%. 8.2.6 Applications Titanium alloy markets and product requirements can be described by three major market segments - jet engines, airframes, and industrial applications, as shown in Table 8-16 (Seagle and Wood, 1993). The first two of these segments are related to the broad aerospace market which dominates the use of titanium in the U.S.A. and consumes about equal amounts for engines and airframes. These two applications are based primarily on the high specific strength (strength-to-density ratio) of titanium. The third, and smallest, market segment in the U.S.A. comprises industrial applications, based on the unique corrosion resistance of titanium towards salt and other aggressive environments. As indicated in Table 8-16, the specified market segments have similar proportions in both the U.S.A. and Europe, although the total U.S. market is about 2.5 times that of Europe based on 1990 data. In Japan, a ma-

431

Table 8-16. Titanium alloys - markets and product requirements. Market segment

Market share

Product requirements

USA Europe Jet engines

42% 37%

Elevated temp, tensile strength, creep strength elevated temp, stability fatigue strength fracture toughness

Airframes

38% 33%

high tensile strength fatigue strength fracture toughness fabricable

Industrial 20% 30%

corrosion resistance adequate strength fabricable cost competitive

Total 1990 consumption (106 kg). USA: 23.6; Europe: 9.1.

jority of the titanium is for non-aerospace use. The titanium capacity of the former Soviet Union was estimated to be about 90 million kg per year; this capacity could totally change the western world marketplace with low cost products. The product requirements for titanium alloys in each market segment are based on the specific needs for the particular application. For example, jet engine requirements are focused primarily on high-temperature tensile and creep strength as well as thermal stability at elevated temperatures. Second tier property considerations are fatigue strength and fracture toughness. Airframe applications require high tensile strength combined with good fatigue strength and fracture toughness. Ease of fabrication of components is also an important consideration. Industrial applications demand good corrosion resistance in a variety of media as a prime consideration as well as adequate strength,

432

8 Titanium, Zirconium, and Hafnium

Figure 8-23. Titanium fan blades for jet engines (courtesy of RMI Titanium Company).

Figure 8-22. Ti-6A1 fan disc forgings for General Electric's CF6 series engine. Each forging is 90 cm in diameter and weighs 250 kg (courtesy of WymanGordon Company).

fabricability and competitive cost, relative to other types of corrosion-resistant alloys. Jet engine applications include discs and fan bladers (Figs. 8-22 and 8-23). Air

frame components produced from titanium vary from small parts to large main landing-gear support beams, the aft section of furlages and truck beam forgings (Figs. 8-24, 8-25, 8-26). Traditional non-aerospace applications cover tubing in heat transfer equipment (Fig. 8-27) and medical prosthesis devices (Fig. 8-28). They also include watches (Fig. 8-29), sporting goods (Fig. 8-30) and

Figure 8-24. Ti-6A1-4V main landing gear support beam forging for the Boeing 747. Each forging is 6.2 m long, 97 cm wide, 28 cm thick and weighs over 1600 kg (courtesy of Wyman-Gordon Company).

8.2 Titanium

433

Figure 8-25. Aft Ti-6A14V/Ti-8Mn "boat-tail" section of fuselage of F-5. This section of the plane experiences heating due to its proximity to the engine (courtesy of Northrop Corporation, Aircraft Division).

Figure 8-26. Boeing 777 Ti-10V-2Fe-3Al truck beam forging; a welded assembly about 10 m long (courtesy of Boeing Commercial Airplane Company).

Figure 8-27. Titanium tubing in heat transfer equipment (courtesy of RMI Titanium Company),

434

8 Titanium, Zirconium, and Hafnium

Figure 8-28. Titanium bar is used for medical prosthesis devices (courtesy of RMI Titanium Company).

8.2.7 Concluding Remarks and Future Thoughts

Titanium possesses a set of novel characteristics including:

Figure 8-29. Titanium watch cases (courtesy of Titan Titanium).

roofs of buildings (Fig. 8-31). Clearly an area for expansion for titanium is in automobiles; an "all titanium" automobile was produced in the mid-1950s (Fig. 8-32). However, wide-spread use for large volume production of automobiles, although contemplated (Fig. 8-33), will require a cost-effective product.

• Titanium is plentiful, being the structural metal which is fourth most abundant in the earth's crust. • Titanium alloys easily with many other elements, which has led to a raft of commercial alloys. • Because of its excellent combination of mechanical properties, low density, and outstanding resistance to hostile environments, titanium is widely used in the aerospace industry for airframes, engines, and rockets. • The very good corrosion resistance of titanium, especially under oxidizing conditions, has led to widespread use in many chemical processing applications and for body implant parts. • Titanium is an attractive metal which can easily be anodized to a variety of colors. This has given rise to applications in buildings and jewelry. • The processing of titanium components is well established in terms of both shape production and development of required microstructures, the latter having

8.2 Titanium

435

Figure 8-30. Experimental high performance "Tiphoon" Ti-15V-3Al-3Cr-3Sn softball bat (courtesy Easton Sports).

Figure 8-31. Seam-welded titanium roof, Futtsu Technology Center, Japan (courtesy of Nippon Steel).

Figure 8-32. All-titanium 1956 GM Titanium Firebird 2 automobile.

Fixture (Door Mjrmx) Valvesprings

Coil Springs

Retainers Front Bumpers & Lowers

Rear Bumpers & towers

Door Panels Door Sill Covers LugNuls Wrmete

Connecting Rods

Valves

Figure 8-33. Prime components for titanium substitution in large production volume automobile (courtesy of Japan Titanium Society).

436

8 Titanium, Zirconium, and Hafnium

a strong influence on mechanical properties. • Titanium is reactive at elevated temperatures, a particular concern being the interaction with oxygen, nitrogen and hydrogen. • Titanium easily forms stable intermetallic compounds. The titanium aluminides (TixAl, where x = l or 3) are nearing commercial applications because of their excellent elevated temperature specific (density normalized) mechanical properties. Lack of ambient temperature "forgiveness" (ductility, fracture toughness, fatigue crack growth rate, etc.) is, however, a great concern. • Titanium is inherently expensive and many approaches to reducing the cost of components - including casting and powder metallurgy methods - have been developed with varying levels of success. Titanium has been referred to as the "wonder" metal. However, to those closely associated with this material, a better word might be the "frustrating" metal. With so much going for it, why is the market in the U.S.A. currently only at the 16 million kg per year level? Over the years, problems such as hydrogen embrittlement and stress corrosion cracking have raised major concerns, with the problems subsequently being overcome. So what is the underlying reason for the market not being much larger? The answer is cost. Cost must be reduced (to perhaps < $7 per kg) to enter the potentially extremely large volume automobile market-place. Much effort is currently directed towards reducing cost, using both near net shape technologies and inherently lower cost material, the latter by relaxing the stringent aerospace chemistry and processing specifications.

8.3 Zirconium 8.3.1 History Klaproth, a German chemist, analyzed a variety of the mineral zircon and discovered zirconium in 1789. However, the metal itself was not isolated until 1824, when Berzelius produced a brittle, impure metal powder by reduction of potassium fluorozirconate with potassium. A hundred years later in Eindhoven, Holland, van Arkel and de Boer developed the iodide decomposition process to make a pure, ductile metal. The "Iodide Crystal Bar" process continues to be used today as a method of purifying titanium, zirconium, and hafnium, even though it is slow and expensive. In the 1940s, several groups of scientists and engineers were investigating zirconium and other metals for nuclear reactors. For this application, a suitable structural metal should exhibit good high-temperature corrosion resistance in water, resistance to irradiation damage, and transparency to thermal neutrons needed to sustain the nuclear reaction. There was a renewed interest in developing a process that could produce a large quantity of zirconium at a much lower cost. In 1945, only a few hundred pounds of zirconium were produced in the U.S. The cost was over $650 per kg. In 1945, development work on zirconium was initiated at the U.S. Bureau of Mines in Albany, Oregon, under Dr. Kroll's technical direction. Already, before 1940, Dr. Kroll had developed a production process for titanium through reduction of titanium tetrachloride with magnesium in an inert atmosphere. A similar process for zirconium was developed in 1947 as a pilot plant with a weekly capacity of 27 kg of zirconium sponge.

8.3 Zirconium

About the time of Kroll's work, Dr. Kaufman of MIT and Dr. Pomerance of Oak Ridge noticed that zirconium, as occurring in nature, was combined with hafnium. It was the hafnium which gave the zirconium the high level of neutron absorption. When the hafnium was removed, zirconium was found to have a very low thermal neutron absorption cross section. This was a scientific fact of great importance. At once, Admiral Rickover, who directed the U. S. Navy Nuclear Propulsion Program, decided to choose zirconium for the naval reactor. This decision stimulated an array of R&D programs to advance zirconium technology in production, Zr/ Hf separation, property information, fabrication, and applications. It was found that highly or commercially pure zirconium was not ideal because of its inconsistent corrosion and oxidation resistance in high-temperature water and steam. This abnormal behavior was attributed to the presence of minor impurities. In particular, the effect of nitrogen upon the corrosion characteristics was very pronounced. Various alloy development programs were established in the early 1950s to examine the effects of adding various elements to zirconium. Independent discoveries by Battelle Memorial Institute and Iowa State College revealed that tin proved most beneficial. The Zr-2.5 %Sn alloy was named Zircaloy-1, which was recommended for the Nautilus reactor. By 1952, data showed that Zircaloy-1 had an increasing rate of corrosion over time and activities on Zircaloy-1 were stopped. An urgent search for a new alloy began. Fortunately, Bettis Atomic Power Lab already had an active program of corrosion tests for a number of zirconium-based alloys. Included was one ingot to which a small amount of stainless steel had been accidentally added. Test results revealed

437

the beneficial effects of iron, nickel and chromium. Quickly, Zircaloy-2, i.e., Z r 1.5%Sn-0.12%Fe-0.1%Cr-0.05%Ni, was developed and specified for the Nautilus reactor in August 1952. The reactor started to generate power on December 30, 1954. The Nautilus got underway on January 17, 1955. These events marked the beginning of a new era. Developmental work continued, since the limiting factor for Zircaloy-2 in a reactor was determined to be its absorption of hydrogen during corrosion in high-temperature water. Bettis eventually discovered that replacing nickel with iron produced an alloy which cut hydrogen absorption by half. This alloy was named Zircaloy-4. There is a controversy on the effect of nickel. Some believe that the nickel addition improves zirconium's corrosion resistance, others don't. Nevertheless, both Zircaloys are important materials for nuclear technology. Demand for zirconium was on the rise, as the U. S. Congress had authorized several nuclear submarines by the mid-1950s, and nuclear power plants were on the horizon. The production cost for zirconium needed to be lowered. This goal could be achieved by developing commercial sources, which included Carborundum Metals, National Distillers Products, NRC Metals, and Wah Chang. Wah Chang contracted to provide zirconium at a price just under $25 per kg in April, 1956. Today Wah Chang, now Teledyne Wah Chang (TWC), remains the most experienced zirconium producer. In 1958, zirconium also became available outside U.S. Navy programs. Activities in developing applications for zirconium were booming. The chemical process industry began to use zirconium by taking advantage of its excellent resistance to a broad range of corrosives. Thanks to its

438

8 Titanium, Zirconium, and Hafnium

remarkable corrosion resistance and biocompatibility, zirconium has found some medical applications, in surgical tools and instruments, and for stitches for brain operations. Zirconium is highly beneficial as an alloying element in iron-, copper-, magnesium-, aluminum-, molybdenum-, and titanium-based alloys. Its ability to combine with gases when hot makes zirconium useful as a getter. Along with niobium, zirconium is superconductive at low temperatures and is used to make superconductive magnets. In the nuclear industry, stainless steel was used to clad the uranium-dioxide fuel in the first generation reactors. But by 1965, the force of neutron economy had made zirconium alloys the predominant cladding material for water-cooled reactors. There was a widespread effort to develop strong, corrosion-resistant zirconium alloys. Worth mentioning, the Ozhennite alloys were developed in the Soviet Union for use in pressurized water and steam. These alloys contain tin, iron, nickel, and niobium, with a total alloy content of 0.51.5%. The Zr-1 %Nb alloy is also used in the Soviet Union for pressurized water and steam service. Researchers at Atomic Energy of Canada Ltd. took a lead from the Russian zirconium-niobium alloys and developed the Zr-2.5%Nb alloy, which is strong and heat treatable. It is used either in a cold-worked or in a quenched-andaged condition. Zirconium is often stated to be a rare, exotic metal. But in fact, zirconium is plentiful and is ranked 19th in abundance of the chemical elements occurring in the earth's crust. It is more abundant than many common metals, such as nickel, copper, chromium, zinc, lead, and cobalt. The most important source for zirconium is zircon (ZrO 2 • SiO2), which appears in several regions throughout the world in the

form of beach sand. The supply of zirconium will not be a problem in developing any application for this metal and its alloys. Moreover, the cost of zirconium has been stable for many years and is competitive with other high-performance materials. Additional general information on zirconium is available in the references given in the "General Reading" section. 8.3.2 General Characteristics

Zirconium alloys are classified in two major categories: nuclear and non-nuclear grades. They all are low in alloy content. They are based on the a structure with dilute additions of solid solution strengthening and a-stabilizing elements such as oxygen and tin. However, in niobium-containing alloys, there are small amounts of niobium-rich /? particles. The most important difference between nuclear and nonnuclear zirconium alloys is in their hafnium content. Nuclear grades of zirconium alloys are essentially hafnium-free ( < 100 ppm). Non-nuclear grades of zirconium alloys may contain up to 4.5% hafnium. Hafnium has a dramatic effect on the nuclear properties of zriconium, but has little effect on its mechanical and chemical properties. The majority of nuclear-grade materials is tubing, which is used for fuel rod claddings, guide tubes, pressure tubes, and ferrule spacer grids. Flat products, such as sheets and plates, are used for spacer grids, water channels and channel boxes for fuel bundles. Bars are used for fuel rod end plugs. A broad range of zirconium products are available for non-nuclear applications. These products include tubes, pipes, sheets, plates, wires, bars, foils, and castings of various sizes. They are used to con-

8.3 Zirconium

struct highly corrosion-resistant equipment, such as heat exchangers, condensers, reactors, piping systems, columns, pumps, valves, and packings for use in the chemical process industries. 8.3.2.1 Physical Properties Zirconium is a lustrous, grayish-white, strong, ductile metal. A summary of physical properties of zirconium is given in Table 8-17. The table indicates that the density of zirconium is considerably lower than those of iron- and nickel-base stainless alloys. In addition, zirconium has a low coefficient of thermal expansion, favoring equipment that requires a close tolerance. The coefficient of thermal expansion of zirconium is about two-thirds that of titanium, about one-third that of type 316 stainless steel (SS), and about one-half that of Monel. Furthermore, zirconium has high thermal conductivity, which is more than 30 % better than in the case of stainless alloys. These properties make zirconium very interesting for constructing compact, efficient equipment, 8.3.2.2 Alloy Behavior Zirconium, like titanium, exists in two crystalline states: the low-temperature a phase, which has a close-packed hexagonal crystal structure, and a high-temperature ft phase, which has a body-centered cubic structure. The allotropic transformation occurs at 865 °C. a-Zirconium is a poor solvent for most elements and practically, /J-zirconium is not obtainable by quenching. The addition of most elements to zirconium results in complicated microstructures which have undesirable properties such as brittleness and degraded corrosion resistance. Only small number of alloying elements can be utilized to improve zirconium's properties.

439

Table 8-17. Physical and mechanical properties of zirconium. Atomic number Atomic weight Atomic radius 0 charge 4+ charge Density Crystal structure a phase P phase Melting point Boiling point Coefficient of thermal expansion Thermal conductivity (300-800 K) Specific heat Vapor pressure at 2273 K at 3873 K Electrical resistivity at 293 K Temperature coefficient of resistivity Latent heat of fusion Latent heat of vaporization Modulus of elasticity Shear modulus Poisson's ratio (ambient temperature)

40 91.22 0.160-0.162 nm 0.080-0.090 nm 6510 kg/m3 hexagonal closepacked <1138K body-centered cubic >1138K 2125 K 4650 K 5.89xlO" 6 K" 1 22 W (m K) 285J(kgK) 0.01 mmHg 900.0 mmHg 0.397 uO m 0.0044 K" 1 253 MJ/kg 6490 MJ/kg 99GPa 36GPa 0.35

Fortunately, zirconium is one of few metals that resists attack by strong acids and alkalies. There are very limited areas for zirconium to improve its corrosion resistance, such as pitting resistance in oxidizing chloride solutions. To develop pittingresistant zirconium alloys continues to be a major research topic. However, for nuclear applications, the oxidation resistance of zirconium is not ideal in high-temperature water or steam. Also, its strength and creep resistance are only marginal at elevated temperatures.

440

8 Titanium, Zirconium, and Hafnium

Alloy development has resulted in some success in improving these properties in order to take advantage of zirconium's transparency to thermal neutrons and resistance to radiation damage. Zirconium exhibits very variable behaviors in hot water and steam. Typically, it will develop black adherent oxide films which remain adherent for a time ranging from short to considerable. After this time, breakaway oxidation can occur. This abnormal behavior was found to be associated with minor impurities such as nitrogen and carbon. The addition of tin and niobium can suppress the pronounced effects created by these impurities. Zirconium alloys, such as Zircaloys and Zr-2.5 wt. %Nb, have been developed for improved, predicable oxidation behavior in hot water and steam. These alloys also have improved mechanical properties.

8.3.2.3 Phase Diagrams In contrast to titanium, zirconium does not alloy easily. a-Zirconium, stable at normal temperatures, is a difficult solvent for most elements, exceptions being titanium, hafnium, scandium, and oxygen. It reacts with most elements to form intermetallics at elevated temperatures. For example, iron is a common ingredient in zirconium, and there are several zirconiumiron compounds, as shown in Fig. 8-34. The very low solubility of iron in a-zirconium is indicated in Fig. 8-35. Although j8-zirconium, stable at temperatures above 865 °C, is a much better solvent than a-zirconium, it is almost impossible to obtain in a metastable /J-state by quenching. Two phase diagrams of engineering importance, i.e., Zr-Sn and Zr-Nb, are shown in Figs. 8.36 and 8.37. Zircaloys

Weight Percent Iron 10

c

20

30

40

50

60

70

80

90 10

2000 j 1855 °C 1800 1600

\

• \ \

1673°C

•

\

v

L \

s'

1400

o o

-^^66.7'i N ^ 1 ^ \ (6Fe) ; Y -482-<^-95~sl^ 1538 °C ! 1337°C V V i - 1394 °C

\

-V
\

x

974 °C

1000 863 °C 800

f2

\4.0 600 400

730 °C

f

775°C

N

;

0.03

i

— (ctZr)

I

MagneticTrans. -300 °C

j

j

-99.9 770 °C ; Magnetic Trans.

1^470 °C («Fe)—— I -275 °C ; Magnetic Trans.

CO J

200 00 Zr

6

1357°C! ~99.3 (yFe)-—I 925 °C 912°C

j | i

1200

I

10

20

30

40 50 60 Atomic Percent Iron

70

Figure 8-34. Assessed Zr-Fe Phase diagram (Arias and Abriata, 1988).

80

90

100 Fe

8.3 Zirconium Weight Percent Iron 0.01 0.02

441

0.03

800 730 °C O

700 (aZr)

o

a

600

o 85StU Mossbauer spectroscopy

500

Figure 8-35. Solubility limit of Fe in a-Zr (Arias and Abriata, 1988).

400 0.01

0.02 0.03 Atomic Percent Iron

0.04

(i.e., 1.5 wt. % Sn-containing alloys) and Zr-2.5 wt. % Nb are the most common commercial alloys. Moreover, Z r - 1 wt. % Nb, which is not included in ASTM specifications, is popular in Russia. Tin stabilizes the hexagonal a phase, and niobium stabilizes the cubic /? phase. Since only small amounts of alloying elements are added to zirconium, all zirconium alloys are basically a alloys. The addition of up to 2.5 wt. % Nb to zirconium may produce some /?-phase particles in the a-phase matrix. 8.3.2.4 Microstructural Development

The close-packed hexagonal lattice of a-Zr entails strong directional variation of mechanical properties of a single crystal. This anisotropy is more noticeable in plastic deformation, which occurs by slip on the prism planes in the crystallographic structure and also by twinning on several different systems. In plastic deformation of a polycrystalline metal it is estimated that about 15 % of the deformation is a result of twinning. Zirconium and its alloys invariably display a strong tendency for a preferred orientation of the crystals,

0.05

known as texture. Texture develops in all stages of production. Well-controlled processes can produce zirconium materials of a favorable texture for strength and resistance to stress corrosion cracking. Cold working on zirconium alloys rapidly increases their hardness, but after 10-15% reduction of area, the rate of hardening is reduced. On heating a coldworked material, the hardness drops. Recovery will take place up to about 500 °C, and somewhere between 500 and 600 °C recrystallization occurs. Zirconium alloys have complex microstructures. In the annealed condition, commercial zirconium alloys have microstructures of equiaxed a-grains. When they are heated into the (a +/J)-phase region and slowly cooled to room temperature, second phase particles precipitate at the grain boundaries of the zircaloys, but the microstructure of annealed Zr-2.5 wt. % Nb remains unchanged. When the alloys are slow-cooled from the j8-phase into the (a + /?)-phase region, a nucleates at the /?-grain boundaries and grows to form a Widmanstatten structure. When Zr-2.5 wt. % Nb is water quenched from the /?- or high (a + /?)-phase regions

442

8 Titanium, Zirconium, and Hafnium

0

10

20

30

40

Weight Percent Tin 50 60 70

80

100

90

2200

200 0 Zr

10

20

30

40 50 60 Atomic Percent Tin

70

80

90

100 Sn

Figure 8-36. Zr-Sn phase diagram (Abriata et al., 1983).

Weight Percent Niobium

-r 1600 (jiZr.Nb)

3 1400

I = 1200 1000 866 °C

977 °C

800 18.7

600

92.1

-(a*) 400 200 0 Zr

10

20

30

40 50 60 Atomic Percent Niobium

Figure 8-37. Z r - N b phase diagram (Okamoto, 1992).

70

80

90

100 Nb

8.3 Zirconium

the j8 transforms to a', which has a martensitic structure. Moreover, a metastable hexagonal co-phase can form in Z r - N b alloys either by quenching the (3 phase or by aging a niobium rich (} phase. The strength of zircaloys cannot be increased very much by heat treatment. The higher strength of the j8 quenched zircaloys compared to the slow-cooled material is partially due to the finer grain size of the Widmanstatten structure. The Zr-2.5 wt. % Nb alloy is much stronger in the /? or (a + /J) quenched condition than in other conditions due to the martensitic a' phase. Aging the a' phase at 500 °C has only a small effect on the strength. / Interstitial impurities, i.e., H, C, N and O, have a pronounced effect on the mechanical properties of zirconium alloys. Carbon, nitrogen and hydrogen are usually kept at a low level because of their harmful effects. Oxygen, on the other hand, is used to strengthen zirconium alloys and is often specified to a level of about 1200 ppm. 8.3.3 Processing/Fabrication Zirconium alloys are ductile and quite workable. Conventional equipment can be used for machining, forming, welding, casting, etc. In certain cases, a few modifications and special techniques are needed. Of primary concern is their tendency to react with gases in air at elevated temperatures and to gall and seize under sliding contact with other metals. Zirconium alloys actually have a better weldability than some more common alloys, provided recommended procedures are followed. For example, zirconium has a low coefficient of thermal expansion which contributes to a low distortion during welding, and no fluxes are needed in welding. One very important aspect in the

443

welding of zirconium alloys is the cleaning and proper shielding of the welding area. 8.3.4 Mechanical Properties Zirconium ores generally contain a few percent of its sister element, hafnium. Hafnium has chemical and metallurgical properties similar to those of zirconium (see later), although their nuclear properties are markedly different. Hafnium is a neutron absorber, but zirconium is not. As a result, there are nuclear and non-nuclear grades of zirconium and zirconium alloys. The nuclear grades are essentially hafnium-free ( < 100 ppm), and the non-nuclear grades may contain up to 4.5% Hf. Some commercially available grades of zirconium and its alloys are listed in Table 8-18. The presence of hafnium in zirconium does not significantly influence mechanical properties other than thermal neutron cross section. Moreover, hafnium is a valuable metal for many applications. Hafnium arises as a by-product in the production of zirconium. The non-nuclear grades of zirconium alloys are also low in hafnium content. Consequently, the counterparts of nuclear and non-nuclear grades of zirconium alloys are interchangeable in mechanical properties. However, specification requirements for nuclear materials are more extensive than those for non-nuclear materials. Only requirements for non-nuclear materials are given in Tables 8-19 and 8-20. It can be seen that Zr 705 is the strongest one with an excellent fabricability. Furthermore, Zr 706 has been developed for severe forming applications by lowering the oxygen content of Zr 705. Additional typical property data are given in Table 8-21 and Figs 8-38 and 8-39.

444

8 Titanium, Zirconium, and Hafnium

Table 8-18. Commercially available grades of zirconium alloys. Alloy design (UNS no.)

Composition (%) Zr + Hf (min)

Ni

Hf (max)

Sn

Nb

Zircaloy-2 (R60802)

0.010

1.20-1.70

-

0.07-0.20 0.05-0.15 0.03-0.08

Zircaloy-4 (R60804)

0.010

1.20-1.70

-

0.18-0.24 0.07-0.13

Zr-2.5Nb (R60901)

0.010

-

2.40-2.80

Fe

Cr

Fe + Cr

Fe-|-Cr O + Ni (max)

Nuclear grades:

-

-

-

0.18-0.38

-

-

-

0.28-0.37 -

-

Chemical grades: Zr702 (R60702)

4.5

0.2 max

0.16

0.2-0.4

0.18

Zr704 (R60704)

97.5

4.5

1.0-2.0

Zr705 (R60705)

95.5

4.5

-

2.0-3.0

0.2 max

0.18

Zr706 (R60706)

95.5

4.5

-

2.0-3.0

0.2 max

0.16

8.3.5 Chemical Properties/Corrosion Behavior

Zirconium is highly reactive, as demonstrated by its redox potential of —1.53 V versus the normal hydrogen electrode at 25 °C. It has a strong affinity for oxygen. In an oxygen-containing medium, zirconium reacts with oxygen at ambient temperature and below to form an adherent, protective oxide film on its surface. This film is self-healing and protects the base metal from chemical and mechanical attack at temperatures up to 300 °C. As a result, zirconium is a highly corrosion-resistant metal. Many engineering metals, such as iron, nickel, chromium, and titanium, produce metal ions of variable valency. Zirconium is predominantly tetravalent in its oxides and many other compounds. It forms very few compounds in which its valence is

other than four. The chemistry of zirconium is characterized by the difficulty of achieving an oxidation state less than four. This characteristic, along with high oxygen affinity, allows zirconium to form protective oxide films even in highly reducing media, such as hydrochloric acid and dilute sulfuric acid. Under these conditions, common metals and alloys may form subordinate oxides or other compounds of low or no protective capability. Moreover, metal ions of constant valency imply stability. This is an important requirement in many applications. For example, ZrCl 4 is used as a catalyst in such reactions as the cracking of petroleum and polymerization of ethylene. Also, zirconium can maintain the stability of certain chemicals, such as hydrogen peroxide. Ions of variable valency are the common decomposition catalysts for hydrogen peroxide.

445

8.3 Zirconium

Table 8-19. Non-nuclear minimum ASTM requirements for the room-temperature mechanical properties of zirconium alloys. Minimum tensile strength

Alloy

Zr702 Zr704 Zr705 Zr706

Minimum elongation

(MPa)

Minimum yield strength (0.2% offset) (MPa)

380 414 552 510

207 240 380 345

16 14 16 20

Bend test radius a

(%) 5T 5T 3T 2.5 T

Bend tests are not applicable to material more than 4.75 mm thick. T is the thickness of the bend test specimen.

Table 8-20. ASME mechanical requirements for Zr702 and Zr705 used for unfired pressure vessels. Material form and conditions Flat-rolled products Seamless tubing Welded tubinga Forgings Bar

ASME Alloy Tensile specifi- grade strength cation number (MPa) SB 551 SB 523 SB 523 SB 493 SB 550

702 705 702 705 702 705 702 705 702 705

359 552 359 552 359 552 359 552 359 552

Minimum yield strength (MPa)

Maximum allowable stress in tension for metal temperature not exceeding: (MPa) 40 °C

95 °C

150°C 205 °C 260 °C 315 °C 370 °C

207 379 207 379 207 379 207 379 207 379

90 138 90 138 77 117 90 138 90 138

76 115 76 115 65 97 76 115 76 115

64 98 64 98 55 83 64 98 64 98

48 86 48 86 41 73 48 86 48 86

42 78 42 78 36 66 42 78 42 78

41 72 41 72 35 59 41 72 41 72

33 62 33 62 28 52 33 62 33 62

a

85% joint efficiency was used to determine the allowable stress value for welded tube. Filler material shall not be used in the manufacture of welded tube.

Table 8-21. The 107-fatigue limits for zirconium alloys. Fatigue limit (MPa)

Iodide Zr Zircaloys or Zr705 (annealed 2 h at 732 °C) Zr-2.5%Nb (aged 4 h a t 566 °C)

Smooth

Notched

145 283

55 55

290

55

Protective oxide films are difficult to form on the surface of zirconium in a few media, such as hydrofluoric acid, concentrated sulfuric acid, and oxidizing chloride solutions. Consequently, zirconium is not suitable or needs protection measures for handling these media. Zirconium forms a visible oxide film in air at about 200 °C. The oxidation rate becomes high enough to produce a loose,

446

8 Titanium, Zirconium, and Hafnium T e m p e r a t u r e , °F 100

200

300

100

400 500 600 700 800

200

(a)

300

400

T e m p e r a t u r e , °C

T e m p e r a t u r e , °F 100

200

300

100

400

200

(b)

500 600 700

300

800

900

400

T e m p e r a t u r e , °C

100 ouu (87) 500 (73)

T e m p e r a t u r e , °F 300 400 500 600 700 800 1 1 i i I I

200

\

900

'

s

\

V Ultim
\ \

i

CO CL

N \

£(44) Yield strength x, 1 (^ 200 (29)

/^ /

-T~ 1

—

Elongation

100 (14) n

-

20

-

10 0

100

(c)

- 30

200

300

T e m p e r a t u r e . °C

400

white scale on zirconium at temperatures above 540 °C. At temperatures above 700 °C, it can absorb oxygen and becomes embrittled after prolonged exposure. Zirconium reacts more slowly with nitrogen than with oxygen. Clean zirconium starts the nitride reaction in ultrapure nitrogen at about 900 °C. Temperatures of 1300°C are needed to fully nitride the metal. The nitriding rate can be enhanced by the presence of oxygen in the nitrogen or on the metal surface. The oxide film on zirconium provides an effective barrier to hydrogen absorption up to 760 °C. In an all-hydrogen atmosphere, hydrogen absorption will begin at 310 °C, provided that the oxide film is intact. Zirconium will ultimately become embrittled by forming zirconium hydrides when the limit of hydrogen solubility is exceeded. Hydrogen can be removed from zirconium by prolonged vacuum annealing at temperatures above 760 °C. Zirconium is stable in NH 3 up to about 1000 °C, in most gases (CO, CO 2 , and SO2) up to about 300° to 400 °C, and in halogens up to about 200 °C. Zirconium resists attack in some molten salts. It is very resistant to corrosion by molten sodium hydroxide to temperatures above 1000 °C. It is also fairly resistant to potassium hydroxide. The oxidation properties of zirconium in nitrate salts are similar to those in air. Zirconium resists some types of molten metals, but the corrosion resistance is affected by trace impurities, such as oxygen, hydrogen, or nitrogen, in the specific molten metal. Zirconium has a corrosion rate of less than 25 jum/y in liquid lead to

Figure 8-38. Tensile properties vs. temperature curves for zirconium alloys: (a) Zr702, (b) Zr704, (c) Zr705.

8.3 Zirconium

447

103 100

I

Roomtemperatur< ^ - 9 5°C| 205 ' C _——-- 315 °C

2.

--—

8 100

^

^H

-

—

•

-

10 6 10"

10" 5

10" 3

10- 4 Strain, %/h

(a)

10

I

0.01

103

I

— -

.-t

—

—

"

—

'—

—

-

—

_

—

-.,

—

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100

—

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~ :

- 100

40 °C

—

-

— •

—

——

-150 °C 260 Dp 425 °C i

—

—

- 10

10 10"6

Figure 8-39. Minimum creep rate vs. stress curves for zirconium alloys, (a) Zr702, (b) Zr705. 10"5

(b)

10"4 Strain, %/h

600 °C, lithium to 800 °C, mercury to 100 °C, and sodium to 600 °C. The molten metals known to attack zirconium are zinc, bismuth, and magnesium. 8.3.6 Applications 8.3.6.1 Nuclear Reactor Applications Characteristics Mainly because of the low absorption of thermal neutrons 2 by pure zirconium (i.e. zirconium free of hafnium) the application of zirconium alloys in nuclear reactors started very early. In fact Zircaloy-2 and Zircaloy-4 and also Zr-2.5Nb (for chemi2 Macroscopic thermal neutron cross-section of Zr is 0.18 barns, of Zircaloy-4 is 0.22 barns, and of Hf is 113 barns (Schemel, 1977).

10- 3

0.01

cal compositions refer to Table 8-18 in Sec. 8.3.2) were used as the basic structural material for fuel elements and in core structures. The Russian-type water cooled reactors (VVER and RMBK) mainly use Zr-lNb. Today the use of Zircaloy-2 and -4 in fuel assemblies (FAs) for light water reactors (LWRs) of the boiling water reactor (BWR) and pressurized weater reactor (PWR) type is most important with respect to the quantity. The operating conditions relevant for the application of zirconium alloys in those reactors are shown in Table 8-22. Zircaloys for nuclear applications are also treated in detail by Lemaignan and Motta in Chapter 7 of Volume 10 B of this Series.

448

8 Titanium, Zirconium, and Hafnium

Table 8-22. Some typical reactor parameters with relevance for zirconium alloy technology.

Cladding temperature System pressure Coolant Water chemistry

a

PWR a

BWRb

290°C-350°e

286°C-290°C

155 bar water 2-4 ppm H 2 (added)

70 bar water + steam 0.2 ppm O 2 (in water) (radiolytic)

PWR: pressurized water reactor; water reactor.

b

BWR: boiling

Components for BWR fuel assemblies that are mainly made out of zircaloys are - cladding tubes for fuel rods, - spacer grids, water rods and water channels for FA structures, - channel boxes for fuel assemblies and for PWRs - cladding tubes for fuel rods, - spacer grids and guide tubes for FA skeletons. Nowadays, the corrosion behavior of zirconium alloy components is essential for the economically effective design and fabrication of these fuel assemblies. The dimensional behavior caused by growth and creep is also important. Furthermore, the chemo-mechanical interaction of the cladding tubes with fission products like iodine or cesium from the irradiated UO 2 fuel must be considered. Additionally, the mechanical properties such as strength and ductility have to be taken into account, but they are not quite as important. As can be expected, there are typical influences of the reactor irradiation, mainly the fast neutron flux and fluence, on all operational behavior parameters. Due to the different water reactor chemistry, the corrosion behavior varies in dif-

ferent reactor types. In the oxidative environment of a BWR core specific precautions have to be taken with respect to the material properties. These precautions avoid a nodular type of corrosion which, connected with a certain amount of crud, in the reactor water, mainly copper oxide, may lead to fuel cladding defects. In a reductive environment as in PWR and VVER cores, uniform corrosion can lead to rather high oxide thicknesses. Under normal circumstances, the growth of oxide is mainly controlled by the temperature at the interface between the cladding tube and the coolant. An "accelerated" (uniform) corrosion has also been recognized which - under certain microstructural conditions of the material is controlled by the flux of the fast neutrons rather then by the temperature. Under these accelerated conditions, the oxide tends to spall off the component whereas normally it sticks very closely to the material. Today an oxide thickness up to about 100 [im is accepted as conservatively safe to avoid fuel rod defects by corrosion. The precise values depend on the actual heat flux on the cladding surface. As the corrosion of zirconium alloys occurs, an uptake of hydrogen is observed due to the corrosion reaction Zr + 2H 2 O

(3)

Traditionally, the amount of hydrogen absorbed by the zirconium alloy was restricted to below 500 ppm. However, recent studies indicate that a hydrogen uptake to about 1000 ppm is acceptable. Irradiation-induced growth is typical in nuclear pile behavior characteristic of zirconium alloys . This growth strongly depends on the microstructural texture of the alloy. Experience (Holzer and Stehle, 1986) has shown length changes of up to

449

8.3 Zirconium

1 % in fuel rods at high burnups (Fig. 8-40). This creep deformation is controlled by the operating temperature and the fluence of the fast neutrons. At lower operating temperatures the influence of irradiation is larger than the impact of temperature. At about 400 °C the influence of the operating temperature dominates. In the past, a considerable number of fuel rod failures were observed due to a pellet-cladding interaction (PCI), which resulted in stress corrosion cracking (SCC) reaction of the zirconium alloy component leading to typical cracks through the cladding tube. This occurred during or after a power ramp due to the radial deformation of the cladding as the result of radial extension of the pellet leads, in combination with the release of fission products like iodine (or cesium). One measure to mitigate this effect was successfully introduced by General Electric: an internal liner of pure zirconium was added to the cladding tube. Other fuel designs also propose alloyed zirconium liner material for the same purpose. Development has not yet progressed to the point where the optimal solution can be defined. Fabrication As a first step in the fabrication of nuclear grade zirconium, hafnium which naturally accompanies the zircon "sand" (ZrO2 • SiO2) has to be removed because it has a very high cross section for the absorption of thermal neutrons (see above). In the western world this hafnium separation is performed either by a "liquid/liquid" extraction, where the different solubility of Zr- and Hf-O(SCN)2 in water and in an organic solvent (MIBK) is used, or by another method based on (fractional) distillation of (ZrCl 4 , HfCl4) from molten salt. In Russia fractional crystal-

20

,30 40 50

-J

Li

I

22

10 5 10 Fast Neutron Fluence (E>0.82 MeV)

0.1

GWd/t(U)

11

21

I

cm" 2

Reactor Q A O D . A F between spacer pos. A F at spacer pos.

0.2 • $ [ c m - 2 s - 1 ] — (E>0.82 MeV) 6-8x1013 8-10x1013

®

5 0.5

1.0

I

e=kxt 0 - 65

2.0

100

200

500 Irradiation Time

1000

2000

EFPD

Figure 8-40. Length change and diameter decrease of Zircaloy-4 cladding tubes on fuel rods with prepressurization (Holzer and Stehle, 1986).

lization of K 2 ZrF 6 and K 2 HfF 6 is used for the separation. In the western world the reduction to the metal state of the now Hf-free Zr is performed by a process named after Dr. Kroll who first developed it for the production of Ti (see Sec. 8.3.1). ZrCl 4 is reduced by Mg to a zirconium "sponge". Russia uses electrochemical reduction for the production of the majority of its zirconium. After these steps the alloying, melting, and hot and cold processing are performed

450

8 Titanium, Zirconium, and Hafnium

in basically the same way everywhere in the world. Melting is conducted using consumable electrodes in a similar manner to the process used for titanium (Sec. 8.2.3.2). The normal practice is triple melting, i.e., two remelting steps are added mainly to get a very high degree of purity with respect to remaining traces of ZrCl 4 . The main mill products from these fabrication steps are:

Fig. 8-41 gives a schematic overview of the different sequences of fabrication steps that lead to the different mill products noted above (Weidinger, 1983). The special processing steps between alloying and the last step leading to a zirconium component are important mainly with regard to the corrosion behavior. Three categories of major influences should be mentioned:

- tubes, used for fuel rod claddings, guide tubes, water rods and sometimes also for ferrule spacer grids in LWR fuel assemblies and for pressure tubes in HWRs and RMBK reactors; - sheets or plates, used for spacer grids, water channels, channel boxes for fuel bundles and other structural parts of all types of water reactors; - bars, mainly used for fuel rod end plugs.

- chemical composition and chemical homogeneity, not only of the major alloying elements but also of some of the impurities, such as carbon and silicon; - precipitates: size, size distribution and chemical composition, Fig. 8-42 (Garzarolli and Stehle, 1987); - the microstructure of the Zr matrix, both grain size and texture (orientation of the grains).

Zirconium sand Zr/Hf-separation Zr O«-reduction I Zirconium-sponge Scrap recycling-

TT

•Alloying-melting Zircaloy-ingot 1 st forging hot rolling

(3 -quenching 2 nd forging

1 st hot rolling

extrusion

2 n d hot rolling

annealing Tube-hollow rocking annealing Cladding tube Guide tubes Fuel rod manuf.

cold rolling 1 u. . annealing ] m u I t l P I e

Sheet •-•Spacer manuf. -•Channel manuf.

I cold rolling

i cold rolling

annealing

annealing

Billets

Wires

multiple

I

I

Bars

I—•Spacer manuf.

Endplug manuf.

— • F u e l rod manuf.

Plate manuf. Fuel rod . \ manuf. Fuel assembly J

Structure manuf.

Figure 8-41. Overview of Zircaloy technology for nuclear applications (Weidinger, 1983).

8.3 Zirconium Relative Corrosion Rate

451

corrosion properties has led to the development of new and refined alloy compositions and new fabrication procedures to make these new alloys. Westinghouse proposed a Z r - l N b - l S n alloy with a very special treatment and called it "ZIRLO". Siemens introduced a "Duplex ELS 0.8 Cladding", in which the outside of the cladding tube is a Zr-0.8 Sn alloy, for improvement of corrosion resistance without degradation in mechanical strength. Other fuel designers will follow soon with further modified or new zirconium alloys. These

10

10

0.02

, ,Q , 0.1

^m

0.8

Average Diameter of Precipitates

Figure 8-42. Corrosion of Zircaloy as a function of the size of intermetallic precipitates (Garzarolli and Stehle, 1987).

The first and the last of these categories also have a strong influence on the dimensional behavior (growth and creep) and on strength and ductility. Specifically the texture significantly affects the SCC susceptibility of the zircaloys, and is therefore important for the PCI behavior of fuel rods under power ramp conditions. The uses of zirconium and hafnium in a nuclear reactor assembly are shown in Fig. 8-43. Sophisticated quality assurance and process and product related quality control procedures were developed several years ago. Some of these need to be revised to comply with recently developed requirements, particularly with regard to enhanced corrosion requirements. The increasing requirements for improved

Figure 8-43. Complementary applications of zirconium and hafnium in a nuclear reactor assembly. (1) Stainless steel fuel assembly guide, (2) stainless steel fuel assembly tip plate, (3) inconel expansion spring, (4) Zirealoy-2 end plug, (5) Zircaloy-4 fuel channel, (6) hafnium control blade, (7) Zircaloy-2 (often with an inner liner of pure zirconium) fuel rod, (8) Zircaloy-4 spacer strip, (9) uranium dioxide fuel pellets.

452

8 Titanium, Zirconium, and Hafnium

developments make life difficult for the producers of zirconium alloy materials because the variety of materials and processes is increasing in an almost stagnant market. 8.3.6.2 Industrial For over 30 years, many diverse applications have been developed for zirconium, its alloys and its compounds. Zircon is an important refractory material for many abrasive and/or high-temperature applications. It is used as the basic mold material in the foundry industry and tundish nozzles. High-performance ceramics can be produced by stabilizing zirconium oxide (zirconia) with yttria or calcia. They can be used for demanding applications, such as the solid electrolyte in oxygen sensors, high-temperature fuel cells, ceramic tubing, extrusion dies, insulating fibers, diamond substitutes in jewelry, etc. The uses of other zirconium compounds result primarily from the ability of zirconium to complex with carboxyl groups and to form insoluble organic compounds. Ammonium zirconium carbonate enhances the fungicidal action of copper salts on cotton cloth. Both ammonium zirconium carbonate and zirconium acetate have been used for water-proofing of fabrics. Freshly prepared zirconium sulfate solutions replace chromium solutions in the tanning of leather. Both zirconium sulfate and potassium hexaftuorozirconate have been used for flame-resistant treatment of wool fabric. Zirconium carbonate is used for paint driers. Certain anti-perspirant products contain aluminum zirconium tetrachlorohydrex glycine. Because of its reactivity, zirconium is used in military ordinance, including percussion-primer compositions, delay fuses,

tracers, and pyrophoric shrapnel, in getters for vacuum tubes and inert-gas glove boxes, and as shredded foil in flashbulbs for photography and excitation of lasers. Alloying applications of zirconium include zirconium-niobium superconductors, titanium or niobium alloys for the aerospace industry, and strengthening of copper alloys. Currently, there are active programs in developing zirconium alloys for rechargeable batteries and hydrogen storage. Increasingly, zirconium and its alloys are being used as structural materials in fabricating columns, reactors, heat exchangers, pumps, piping systems, valves, and agitators for the chemical process industry (CPI). This will be discussed to a greater extent in the following sections. Zirconium was found to be corrosionresistant in strong acids and alkalies by Gillett in 1940. This resistance was confirmed by Dr. Kroll in 1946 after he developed the commercial production process for zirconium. In this respect, zirconium beats other high-performance materials such as titanium, tantalum, glass, and polytetrafluoroethylene, which are attacked by strong alkalies. This early work, along with later ones, has led to the use of zirconium in solving corrosion problems in modern processes since the 1950s. Modern processes emphasize low costs (raw materials, operation, maintenance, etc.), improved efficiency, high quality, variety and choice, safety, and environmental protection. These requirements often necessitate that modern processes take place at elevated temperatures/pressures in the presence of a catalyst. Therefore, as the CPI modernizes its technologies, process environment may become more corrosive and require more corrosion-resistant equipment to cope with these problems. Corrosion problems continue to mean that process equipment suffers visible, ex-

453

8.3 Zirconium 300

cessive corrosion. Hence, process equipment becomes unsafe and requires repair or replacement. Today, corrosion problems have a much broader meaning because of increasing concerns for efficiency, quality, and environment. A corrosion problem may exist even if process equipment does not show any visible deterioration. For example, type 316 L stainless steel equipment corrodes at 50 jim/y, which is considered to be low. Still, each day for every 1000 m 2 , about 700 g of iron, 175 g of chromium, 110 g of nickel and 20 g of molybdenum are released into the product, the process medium, or the environment. These levels of discharging, although low, can be undesirable and unacceptable. They reduce the quality of fine chemicals, drugs, foods, and fertilizers. They can poison certain catalysts to reduce process efficiency, and cause certain chemicals, such as hydrogen peroxide, to catalytically decompose. In addition, they may be harmful to the environment. Zirconium is not just corrosion-resistant but also nontoxic. It is compatible with many corrosives. Mineral acids are among the most corrosive media for common metals and alloys. The corrosion resistance of zirconium in four mineral acids is illustrated in Figs. 8-44 to 8-47. The CPI has always recognized the advantages of zirconium for solving corrosion problems. Dr. Kroll predicted that zirconium would be useful in HC1 applications, since HC1 is one of the most difficult acids for common metals to handle. Many R&D programs were established to evaluate zirconium for applications that involved harsh conditions (Gee et al., 1949; Golden et al., 1952, 1953; Lane et al., 1953; Gegner and Wilson, 1950). It has proved to be one of the most corrosion-re-

220

>5 mni/yr

5

-0.5

260

°A\

, _

^ .__,

500

-

400

V_

0.05-5 mm/yr \ .13-0.5

180

=

t/yr j

140

y

0-0.13 mm/yr |^—-^ Boilin g point :urve

100

P\\

300 g. E

- 200

60 - 100 20 0

20

40

60

80

100

C o n c e n t r a t i o n of H 2 SO 4 , %

Figure 8-44. The isocorrosion curve of zirconium in sulfuric acid.

I Q.

E

0

10

20

30

40

C o n c e n t r a t i o n of HCI, %

Figure 8-45. The iso-corrosion curve of zirconium in hydrochloric acid.

sistant materials available to the CPI (Yau and Webster, 1987). However, zirconium is more than just that. It is also lightweight, has good thermal conductivity and adequate strength, and it is nontoxic and biocompatible.

454

8 Titanium, Zirconium, and Hafnium

260 i

220

I i

|No1t tested - 350 I I I 1

180 0-0 13 mm) yr

2

- 450

I

140

1

CD Q.

ECD

H

I

.—-—

—'— - — ' Boilin g point curve

100

- 250

\ - 150

60 0-C .13 mm/yr

20 20

40

60

8

100

C o n c e n t r a t i o n of H N O 3 , '

Figure 8-46. The iso-corrosion curve of zirconium in nitric acid.

0

20

40

60

80

100

C o n c e n t r a t i o n of H 3 PO 4 , %

Figure 8-47. The iso-corrosion curve of zirconium in phosphoric acid.

Thus, it is not just used to solve corrosion problems, but also to cope with current social/economic challenges. One of the earliest applications of zirconium was in the field of urea production. Certain zirconium vessels and heat exchangers have been in service for over 30 years and still show no signs of corro-

sion. In modern processes, urea is produced by the reaction of NH 3 with CO 2 to form ammonium carbamate followed by dehydration. In order to have a high conversion rate, reactions have to take place at elevated temperatures and pressures. The reactants, particularly the carbamate solution, are too corrosive for stainless alloys unless oxygen is ejected carefully for passivation. Indeed, oxygen injection has been popular in urea plants for corrosion control of stainless-steel equipment for some years. However, oxygen injection has certain drawbacks: - higher costs resulting from the addition of pumps; - lower plant efficiency; - greater safety problems owing to the risk of formation of explosive gas mixtures ; - poor corrosion resistance of stainless steel when oxygen ejection is interrupted. The safety concern is real, since the explosion of stainless-steel equipment has occurred several times. Moreover, there is an increasing concern for the presence of heavy metal ions in fertilizers. Interest in using zirconium in urea production has resumed in recent years. For example, stainless steel tubes with zirconium lining have been developed for carbamate decomposers and/or condensers. Zirconium is an important material for the production of acetic acid by the synthesis of methanol and monoxide. This technology has been studied for over 40 years. However, the process was only commercialized in the 1970s when the corrosion problems of structural materials became manageable. In this technology, the reaction must proceed at high temperatures (150-200°C) and at high pressure in the presence of an iodide as the catalyst. The

8.3 Zirconium

crude acid produced is first separated from the catalyst and then dehydrated and purified in an azeotropic distillation column. The final product is highly pure acetic acid, allowing it to be used in food and pharmaceutical applications. Here, all the factors inducing corrosion problems with stainless alloys are present. The process equipment must be made of the most corrosion-resistant materials available, such as zirconium and its alloys. Zr 7O2 and Zr 7O5 are standard materials for the construction of critical pieces of equipment for acetic acid service and, in recent years, zirconium has been replacing stainless alloys after their failures in acetic acid service. Formic acid production by the hydrolysis of methyl formate, such as the Leonard/ Kemira process, is another modern process that requires zirconium. Formic acid is more highly ionized and therefore more corrosive than acetic acid. Stainless steels can be seriously attacked by intermediate strengths of hot formic acid. Nickel-based alloys may corrode at high rates in the presence of certain impurities and under heat transfer conditions. The performance of titanium in formic acid may be affected by minor factors, such as aeration. In the Leonard/Kemira process, CO gas contacts methanol in the presence of a catalyst to form methyl formate in a reactor. The methyl formate is hydrolyzed in the presence of a catalyst to yield formic acid and methanol, which are then separated by distillation, and the methanol is recycled. Factors such as elevated temperatures and pressures and the presence of water and catalyst make common materials, including glass lining, resin and plastic coatings, and stainless alloys, inadequate as structural materials for this process. Zirconium proves to be the most economical structural material for use in the main equipment for this process.

455

Zirconium has been used in the manufacture of H 2 O 2 . This production process involves the electrolysis of acid sulfates and is very corrosive. At one time, graphite equipment was standard for this process. In the FMC plant in Vancouver, Washington, it was found that zirconium was superior to graphite. FMC used zirconium equipment to produce up to 90% H 2 O 2 . The average maintenance-free life of the heat exchanger was 10 years. Graphite exchangers had failed after 12-18 months of service. The graphite equipment failure was attributed to the leaching of the binder from graphite by the 35% H 2 SO 4 feed, which created a porous condition and ultimately caused failure. Hydrogen peroxide is becoming a preferred oxidant because it is environmentally safe. It has made inroads to many industrial applications, and its consumption is increasing at a rapid pace. It is, however, an unstable chemical. It can be catalytically decomposed by many heavy metal ions. The decomposition reaction is wasteful and may trigger off fire or explosion. Certain peroxide solutions are also corrosive. Zirconium is one of very few metals that is highly resistant to a broad range of peroxide solutions. It does not produce active ions to catalytically decompose peroxide. Thus, it is attractive for many peroxide applications, such as production and handling of fine H 2 O 2 , pulp bleaching, waste treatment, etc. The experience of zirconium in peroxide production led FMC to replace the graphite heat exchangers with zirconium shell and tube exchangers used in the manufacture of acrylic films and fibers. In this application, the H 2 SO 4 concentration was as high as 60% at 150°C. Another major application involving H 2 SO 4 concerns the manufacture of methyl methacrylate and methacrylic acid.

456

8 Titanium, Zirconium, and Hafnium

The system at Rohm and Haas' plant in Deer Park, Texas, includes pressure vessels, columns, heat exchangers, piping systems, pumps and valves made from zirconium. A zirconium unit, which was built more than 20 years ago, is still in service. Zirconium is also widely used for column internals and reboilers in the manufacture of butyl alcohol. The operating conditions are 60 to 65% H 2 SO 4 at temperatures to boiling and slightly above. Zirconium may corrode under upset conditions of elevated concentrations and when impurities such as Fe 3 + are present. However, it has been used in H 2 SO 4 recovery and recycling systems in which fluorides are not present and the acid concentration does not exceed 65%. A major application for zirconium is in iron and steel pickling, using hot 5 to 40% H 2 SO 4 . Further applications include the production of concentrated HC1 and of polymers involving HC1. Zirconium heat exchangers, pumps, and agitators have been used for over 15 years for azo dye coupling reactions. In addition to being very corrosion resistant in this medium, zirconium does not plate out undesirable salts that would change the color and stability of the dyes. Other applications in HC1 include the break-down of cellulose in the food industry and the polymerization of ethylene chloride, which is carried out in HC1 and chlorinated solvents. Zirconium and its alloys have been identified to offer the best prospects from a cost standpoint as materials for HI decomposers in hydrogen production. They resist attack by HI media (gas or liquid) from room temperature to 300 °C. Most stainless alloys have adequate corrosion resistance only at low temperatures. There is an increasing interest in the use of zirconium for HNO 3 service. For example, because of the high degree of concern

over safety, zirconium is chosen as the major structural material for the critical equipment used to reprocess spent nuclear fuels. In most processes involving HNO 3 , stainless steel has been the workhorse for decades. The excellent compatibility between zirconium and HNO 3 was not considered necessary. The situation changed when nitric acid producers started to modernize their technology in the late 1970s. Conventionally, HNO 3 is manufactured by oxidation of ammonia with air over platinum catalysts. The resulting nitric oxide is further oxidized into nitrogen dioxide and then absorbed in water to form HNO 3 . Acid of up to 65% concentration is produced by this process. Higher concentrations are achieved by distilling the dilute acid with a dehydrating agent. Before the 1970s, dual-pressure processes were the dominant means of H N 0 3 production. A typical dual-pressure process operates the converter at about 500 MPa and the absorber at about 1100 MPa. In the late 1970s, Weatherly Inc. introduced a mono-pressure process which operates at 1300-1500 MPa. The advantages of this new process are - greater productivity due to higher operating pressure; - smaller equipment resulting in a lower capital cost; - higher energy recovery capabilities. The first application of this new process came in 1979 when Mississippi Chemical in Yazoo City retrofit their existing plant with a new compressor system to increase pressure for greater productivity and energy efficiency. It was at this point that severe corrosion problems were discovered. Prior to the upgrade, the cooler condenser was constructed of type 304 L stain-

8.3 Zirconium

less-steel tubesheets and type 329 stainlesssteel tubes. Under the previous operating conditions, the cooler condenser exhibited some corrosion, which was managed by plugging tubes and replacing the unit every three to four years. Shortly after the upgrade, with an operating temperature and pressure of 200°C and 1035 MPa, 10% of the type 329 stainless steel tubes were found to be leaking. So the condenser was replaced with a unit using type 310 L stainless steel, which had to be exchanged after 13 months of operation. The original condenser with new tubes of improved grade 329 stainless steel replaced the 310 L unit. Mississippi Chemical began looking for alternatives. In an attempt to find a solution to this problem, TWC conducted autoclave tests on many newer types of stainless steels and zirconium in solutions up to 204 °C and at concentrations up to 65%. Clearly, zirconium was the only suitable material for the mono-pressure process. Corrosion rates of zirconium coupons were consistently below 25 [im/y. The next step was to test zirconium tubes in service. Several tubes were installed into a rebuilt stainless steel condenser. They were destructively examined after 13 months. There were no signs of corrosion. Zirconium tubes were then placed in another condenser for one year. Once again, there were no signs of corrosion. Consequently, Mississippi Chemical replaced its stainless steel cooler condenser with one constructed from zirconium tubes and zirconium/304 L stainless steel explosion-bonded tubesheets. This unit contains over 18 km of zirconium tubing. The zirconium unit has been in service since 1984 and has already outperformed the stainless steel predecessors. Thereafter, several zirconium cooler condensers have been built for other HNO 3 producers.

457

Mono-pressure plants are not the only ones that use zirconium to solve corrosion problems. Certain plants use a distillation process to increase the acid flow through a reboiler and entering a distillation column to drive off water and concentrate the acid. In 1982, Union Chemicals replaced the bottom portion of each of two distillation columns and the tube bundles of each of two reboilers. The lower parts of the columns had been constructed originally with type 304 L stainless steel and exhibited corrosion problems. Titanium was tried, but also failed. While glass-lined steel did not show the corrosion problems experienced by type 304 L stainless steel and titanium, the maintenance costs were found to be unacceptable. Zirconium provides significantly improved corrosion resistance without adding maintenance costs. It also solved the corrosion problem in the reboilers. Prior to the installation of zirconium tube bundles, both 304 L and titanium tube bundles had failed within less than 18 months of operation. With proper design and fabrication, the susceptibility of zirconium to SCC can be suppressed in highly concentrated HNO 3 . For example, an Israeli chemical plant uses zirconium tubes in a U-tube cooler that processes bleached H N 0 3 at concentrations between 98.5 and 99%. The unit cools the acid from 70-75°C to 35-40°C. Previously, U-tube coolers were made from aluminum, which failed within 2 to 12 weeks. The zirconium unit has been in service for over two years, operating 24 hours a day, six days a week. A unique application for zirconium is in processes that cycle between HC1 or H 2 SO 4 and alkaline solutions. One company replaced a lead and brick-lined carbon steel reactor vessel with zirconium because the reaction alternated between hot H 2 SO 4 and caustic. The vessel has been in

458

8 Titanium, Zirconium, and Hafnium

use for several years with no corrosion problems. A zirconium distillation column proved to be suitable and economical in a chlorinated hydrocarbon environment at more than 150°C. The column was built as a replacement for an old brick-lined column. Corrosion of the old column necessitated frequent renovations, with resultant high maintenance costs and plant down time. In the zirconium extraction process, several corrosive chemicals are used, including methyl isobutyl ketone, HC1, ammonium thiocyanate, H 2 SO 4 , and zirconyl chloride. Zirconium has been found to withstand the rigors of this manufacturing process, as demonstrated by the following five examples. • The pumps used to transport the process chemicals are primary candidates for failure. Various materials, including stainless alloys, plastics, and high-silicon cast iron had been tried, but each showed some limitation during this severe service. All wetted pump parts were converted to zirconium. The oldest one has been operating for more than five years and shows no sign of corrosion. • For heat exchanger service used for the same process chemicals, zirconium tubes replaced impervious graphite tubes that were prone to failure. The oldest zirconium heat exchanger has been in service for more than seven years and shows no signs of corrosion. • Steam stripper columns were converted from a furan resin to zirconium. At the 105 °C operating temperature of the columns, the plastic would embrittle and crack. Failure of these columns often resulted in expensive spills. By contrast, zirconium columns require no maintenance due to material failure.

• Zirconium is used in an electrostatic precipitator installed to scrub corrosive gases emitted from rotary kilns. The emitting electrodes and all fixtures exposed to the ammonium-sulfate and chloride-gas process stream were constructed from zirconium. The emitting electrodes are charged with 45 000 V. • Several parts of a crude chlorination scrubber were converted from plastic to zirconium. The plastic parts were replaced approximately every six months. The zirconium replacement is expected to last about 25 years. Zirconium is nontoxic and biocompatible as mentioned before. With its low corrosion rates in many media, it is also ideal for making equipment used in food processing, the manufacture of fine chemicals, and pharmaceutical preparations. Zirconium and its alloys are even suitable for certain implant applications. Other applications for zirconium include heat exchangers exposed to hot seawater, cladding for high-strength or lightweight alloys, and reclaiming valuable chemicals from wastes. Zirconium is an excellent material for laboratory equipment, such as crucibles and autoclaves. Zirconium crucibles have replaced platinum crucibles for handling certain molten salts. Properly oxidized zirconium and zircaloys can be used as insulating washers and measuring devices exposed to high-temperature corrosives. 8.3.7 Concluding Remarks and Future Thoughts

Zirconium possesses a set of novel characteristics including: - It is highly transparent to thermal neutrons. - Zirconium is one of very few metals that

8.4 Hafnium

-

-

resists attack by strong acids and caustics, as well as many salt solutions and molten salts. It is nontoxic and biocompatible. It has adequate strength. Zirconium has a relatively low density, high thermal conductivity, and low coefficient of thermal expansion. Zirconium can be fabricated into almost any shape by conventional methods. It is reactive when its surface area is large. It forms stable intermetallic compounds with many elements and reacts with carbon, nitrogen, etc.

As a result of these characteristics, many nuclear and industrial applications have been developed for zirconium and its alloys. These applications include fuel cladding and pressure tubes for nuclear reactors, process equipment for the CPI, superconducting materials, battery alloys, hydrogen storage alloys, ordnance applications, implant materials, and consumer goods. However, there is still a perception that zirconium is exotic and costly; as mentioned in Sec. 8.3.1, zirconium in fact is plentiful. In the earth's crust, it is more abundant than many common elements, such as nickel, copper, chromium, zinc, lead, and cobalt. The price of zirconium and its compounds has been relatively stable for many years, and it is very competitive with other high-performance materials. There are ample opportunities for zirconium and its chemicals to grow in the coming years. However, for the zirconium market to extend, emphasis must be on developing its non-nuclear applications. Except for a few isolated cases, the nuclear industry does not present a growth industry for zirconium in the foreseeable future.

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8.4 Hafnium 8.4.1 History In 1911, G. Urbain announced the discovery of a new element in some residues remaining after the separation of the lutetium-ytterbium fractions of the rare earths. He named it celtium. However, no one could produce this element in sufficient quantity and purity to verify its existence. The credit of discovering hafnium went to D. Coster and G. von Hevesy in 1923, when they detected this element in zirconium minerals. Hafnium was named to honor the city, Hafnia (an ancient name for Copenhagen), where it was discovered. Hafnium is always found in zirconium minerals, which contain 1-6% hafnium. There is only one mineral found in Scandinavia, thortveitite, which contains more hafnium than zirconium. Zircon and baddeleyite are two important minerals for the production of these elements. The existence of hafnium was hidden behind zirconium for over 130 years. The remarkable similarity between hafnium and zirconium in large part accounts for its late discovery. Of all the elements, zirconium and hafnium are two of the most difficult to separate. Literally hundreds of techniques have been proposed and studied for their separation. Two commercially important techniques used in the western world are the liquid-liquid extraction of hafnium from hydrochloric acid solution by a methyl isobutyl ketone-thiocyanic acid solution (Stickney, 1959) and the extractive distillation of hafnium tetrachloride from the mixed tetrachlorides dissolved in a molten halide solvent (Besson et al., 1976). The separation of zirconium and hafnium was the key in developing nuclear applications for these metals. Regardless of the unusual similarity of their chemical

460

8 Titanium, Zirconium, and Hafnium

properties, zirconium and hafnium are strikingly different in the cross-section absorption for thermal neutrons. Hafnium absorbs thermal neutrons almost 600 times more than zirconium. Consequently, the use of zirconium as a structural material in nuclear reactor cores demands low hafnium contents (below 100 ppm). On the other hand, hafnium is useful as a nuclear fission control-rod material. 8.4.2 General Characteristics Hafnium has been used primarily in its elemental form. There is limited information for hafnium-based alloys (Nielson, 1980), such as Hf-Ti and Hf-Zr alloys, which are not commercially produced. Nevertheless, they all are based on the a structure with dilute additions of solid solution strengthening agents such as oxygen and zirconium. Hafnium is produced as a by-product in zirconium production, and is therefore not available in large quantities. This has limited the chance for hafnium to be considered in applications that require large amounts. However, there are niche applications that do not require high volumes. These applications include alloying, stereo headphones, and active tips for plasma arc cutting tools. Commercially, hafnium is produced as sponge, powder, plates, rods, wires and foils. 8.4.2.1 Physical Properties Hafnium is a heavy, ductile metal with a brilliant silvery luster. Although chemically very similar to zirconium, hafnium has several differences in physical properties : hafnium has twice the density of zirconium, a higher phase transition temperature, a higher melting point, and a much higher thermal neutron absorption coeffl-

Table 8-23. Physical properties of hafnium. Atomic number Relative atomic mass Ar Melting point Boiling point Density Thermal conductivity (298 K) Coefficient of linear expansion (273-1273 K) Specific heat (298 K) Vapor pressure (2040 K) (2280 K) Electrical resistivity (298 K) Thermal neutron absorption cross-section Crystal structure; <x form f$ form Temperature of a-/? transformation

72 178.49 2500 K 4875 K 13 310kg/m 3 23.0 W/(m K) 5.9 x l O " ^ " 1 117J/(kgK) 10" 5 Pa 10~ 4 Pa 0.351 \\Q m 10400 fm2 (104 barns) hexagonal close packed (h.c.p.) body-centered cubic (b.c.c.) 2033 K

cient. Some physical properties of hafnium are given in Table 8-23. 8.4.2.2 Alloy Behavior Only a very limited number of hafniumbased alloys have been developed. Hafnium, with 20-27% tantalum and up to 2% molybdenum, provides excellent oxidation resistance. 8.4.3 Processing/Fabrication Hafnium, with low impurity contents, e.g. less than 500 ppm oxygen, is ductile and quite workable. It can be processed by the same methods as for zirconium. However, the transformation temperature from the a phase to the /? phase is very high at 1760°C. Consequently, for hafnium, all processing operations are carried out in a single a phase range. Furthermore, at a given temperature, the strength of the hex-

8.4 Hafnium

agonal close-packed structure (a phase) is higer than that of the body-centered cubic structure ([! phase). There is the need to process hafnium at higher temperatures than those necessary for zirconium. Typical forging and hot rolling temperatures for hafnium are 1100°C and 900 °C, respectively. The workability is markedly reduced when the oxygen content increases from 500 to 1000 ppm. A higer rolling temperature at 1100°C would be more appropriate for higher oxygen contents. Although hafnium oxidizes less rapidly than titanium or zirconium, it is necessary to minimize the temperature and time of heating thick stock (> 1.27 mm thick) in air. The highly tenacious oxide scale must be removed by machining, grinding, peening, or pickling in a nitric-hydrofluoric acid bath. The cold and warm workability of hafnium is noticeably poorer than that of zirconium. Hafnium should be cold rolled in a series of light passes (5 % to 7 % reduction) to avoid intergranular herring-bone cracks. Annealing is performed when the total strain reaches 25% - 3 5 % . Annealing of thin stock is generally carried out under vacuum or in an inert-gas atmosphere to avoid contamination. The recrystallization range for hafnium is 700800 °C. For short periods of heat-treatment, a stress relief is obtained by treatment at 845 °C and full annealing is effected by treatment at 925 °C. 8.4.4 Mechanical Properties

Hafnium is a relatively hard metal. Annealing becomes necessary when hafnium is cold worked about 25-35%. Ductility suitable for metal working requires a low oxygen content on the order of 500 ppm, about one-quarter of that amount which would render zirconium difficult to fabri-

461

cate. Typical mechanical properties for hafnium are given in Table 8-24. 8.4.5 Chemical Properties/Corrosion Behavior

The ionic radii of hafnium and zirconium are almost identical, resulting from the lanthanide contraction. Both elements exhibit a valence of four. Therefore, the chemistry of hafnium is similar to that of zirconium. Hafnium, like zirconium, is one of the few metals that resist attack by strong acids and alkalies. It is, however, superior to zirconium and zircaloys in corrosion resistance in water and steam, molten alkali metals, and in air. For example, in one of the standard corrosion tests (360 °C water at 3.9 MPa), the weight gain of zircaloy is about 4.5 times higher than that of hafnium. One major difference in corrosion resistance between hafnium and other engineering metals is found in alkaline solutions. Thermodynamic data indicate that stable hafnium oxide can be formed on the metal's surface in alkaline water (Ayer, et al., 1987). Most engineering metals rely on metastable oxides or the presence of an inhibitor for their resistance in alkaline solutions. Consequently, hafnium could prove superior to most engineering metals in alkaline solutions. Indeed, hafnium has been found suitable for handling hot leaching solutions containing caustic soda and sodium peroxide. In aqueous solutions of its compounds, hafnium is always tetravalent. In dilute acids, it slowly hydrolyzes and polymerizes. Hydrated hafnium oxide precipitates from aqueous solutions at about pH 2. The ammonium-hafnium carbonate and potassium-hafnium carbonate complexes are the few inorganic compounds that have

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8 Titanium, Zirconium, and Hafnium

Table 8-24. Typical mechanical properties for fully annealed products of hafnium. Test direction

Ultimate tensile strength (MPa)

Yield strength (0.2% off-set) (MPa)

Elongation

Rod R.T. 315°C

longitudinal longitudinal

480 310

240 125

25 40

Plate R.T. R.T. 315°C 315°C

longitudinal transverse longitudinal transverse

470 450 275 235

195 310 125 165

25 25 45 48

Strip R.T. R.T. 315°C 315°C

longitudinal transverse longitudinal transverse

450 450 275 240

170 275 95 165

30 30 45 50

General properties

Young's modulus (R.T.) Impact strength (subsize izod, V-notch) (R.T.) Creep strength (longitudinal) (Minimum rate 10~6/h at 315°C±55°C) Fatigue strength at 2 x 107 cycles Unnotched (R.T) Notched (R.T.) Unnotched (370 °C) Notched (370 °C)

significant solubility in neutral or slightly basic aqueous solutions. At high temperatures, hafnium has a strong affinity for hydrogen, oxygen, carbon, and nitrogen. Hafnium reacts with hydrogen slowly at about 250 °C and rapidly at about 700 °C; with air or oxygen at about 400 °C; with carbon at about 500 °C; and with nitrogen at about 900 °C. The formed hydrides, oxides, carbides, and nitrides lead to reduced ductility levels. Actually, these compounds are interesting and useful. For example, hydrogen absorption is a reversible process. The process of hydride formation, crushing, and

135 GPa 0.37 kg m 160 MPa 190 MPa 135 MPa 120 MPa 85 MPa

dehydriding is used to convert hafnium metal pieces into powder with little contamination. Hafnium dioxide, nitride, and carbide are stable, refractory compounds. In fact, hafnium carbide is the most heatresistant binary compound known. The melting points of several refractory metals and their compounds are given in Table 8-25. These properties allow hafnium to find many unique applications in various industries. In molten salts, hafnium is also normally tetravalent. However, in anhydrous molten halide salts, Hf4+ can be reduced toHf 3 + andHf 2 + .

8.4 Hafnium

Table 8-25. Melting points of selected refractory metals and their compounds. Material

Melting point (°C)

W Ta Nb Zr Hf

3410 2996 2468 1845 2222

WO 3 Ta 2 O 5 Nb 2 O 5 ZrO 2 HfO2

1473 1800 1780 2715 2812

WN 2 TaN NbN ZrN HgN

we

TaC NbC ZrC HfC

>400 (in vacuum) 3360 2573 2980 3305 2870 3880 3500 3540 3890

8.4.6 Applications 8.4.6.1 Nuclear The property of hafnium that makes it desirable for use in nuclear reactors is its opaqueness to the passage of thermal neutrons. Although the thermal neutron absorption cross section of hafnium (104 barns) is not as high as that of cadmium (2460 barns) or gadolinium (48 000 barns), hafnium has a unique combination of properties, including excellent corrosion resistance, good strength, ease of fabrication and weldability, resistance to irradiation damage, and superb dimensional stability. These properties fit the design needs as absorbers for control rods in naval reactors. In addition, hafnium maintains its efficiency as a neutron absorber even after prolonged exposure to radiation. Consequently, hafnium is a unique control material for light water-cooled reactors.

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Unalloyed hafnium is used in naval reactors, although controlled amounts of oxygen and iron are needed to assure that strength levels are met. Hafnium control rods were used in several early power reactors, such as the Yankee Rowe and Indian Point reactors, for electricity generating plants. However, this commercial application did not continue, due to a perceived limited availability and high cost of hafnium. Boron carbide was chosen for boiling water reactors, and the silver-indium-cadmium alloy was selected for pressurized water reactors. Both materials have since encountered problems. Boron carbide can be easily damaged by irradiation due to swelling. Costs for the manufacture of silver-indium-cadmium rods have escalated faster than for hafnium. Also, silver is a speculative metal. Over the last decade, the cost of hafnium has remained stable since its availability exceeds market demands. Annual western hemisphere availability is estimated at 801, with consumption assessed at 50 t (Nielsen, 1980). As a result, hafnium clad in stainless steel is now replacing boron carbide and the silver-indium-cadmium alloy in some commercial nuclear power plants. In some cases, both hafnium and boron carbide are used to construct control blades for boiling water reactors, as shown in Fig. 8-48 (Morlin and Nordlof, 1985). Furthermore, a new type of hafnium absorber rodlets has been developed for pressurized water reactors (Keller et al., 1982). A unique feature of the design is the sealing of the hafnium material inside a stainless steel tubing. Previously, the hafnium material was exposed directly to the reactor primary coolant. The use of hafnium and zirconium in a nuclear reactor assembly are shown in Fig. 8-43.

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8 Titanium, Zirconium, and Hafnium

Hi-Tip

Hybridzone

B4Czone

Figure 8-48. A control blade for a boiling water reactor (courtesy of ASEA-ATOM).

8.4.6.2 Industrial Hafnium is a highly effective element for the enhancement of mechanical and corrosion properties in a wide variety of alloy systems. It is one of the most effective solid-solution strengtheners. Its affinity for oxygen, nitrogen, and carbon results in precipitation strengthening by formation of second-phase particles. Hafnium plays an important role in optimization of alloy design for strength, ductility, toughness, creep resistance, oxidation resistance, stress corrosion cracking resistance, and fabricability. The largest volume alloying application for hafnium is its addition (1 - 2 %) to nickel-based superalloys. The hafnium content significantly improves performance and

ductility of these alloys. It has raised the allowable service temperature by 50 °C, which results in improved engine efficiency. These alloys are currently used in gas turbines. These alloys are employed in the hottest stages following the combustion chamber, jet engine vanes and rotating blades for both commercial and military aircraft. Examples are TRW VIA (Ni- 6Cr-7.5Co-2Mo-5.8W-l T i 5.4Al-0.13Zr-0.5Nb-9Ta-0.5Re-0.5 Hf), Unitemp C-300 ( N i - 8 . 7 C r - 9 C o 2Mo-7.6W-0.7Ti-3.4Al-10Ta-0.5VlHf), MM-002 ( N i - 9 C r - 1 0 C o - 1 0 M o 1.5Ti-5.5Al-lFe-2.5Ta-1.5Hf), and MM-008 (Ni-14.6Cr-15.2Co-4.4Mo3.35Ti-4.3Al-1.3Hf). Hafnium has been used extensively in tantalum-based alloys to increase the hotstrength. Alloys such as T-lll ( T a - 8 W 2Hf), T-222 (Ta-9.6W-2.4Hf) and Astar 811 C (Ta-8 W - l R e - 1 Hf) have applications in re-entry vehicles, honeycomb heat shields, and aerospace fasteners. Certain niobium-based alloys contain 10% or more hafnium. Alloys C-103 ( N b lOHf-lTi) and C-129Y ( N b - l O W 10Hf-0.07Y) have found widespread application in jet engines and missile systems. These alloys show high hot-strength, good weldability and formability necessary in gas turbine and aerospace applications. Alloy C-103 was used extensively in the Apollo command module engine and is the nozzle alloy on most of the telecommunication satellites. It is also used on the thruster nozzle of space shuttles. The largest use for Alloy C-103 is in thrust augmenters for jet engines. Another alloy, WC-3015 (Nb-30Hf-15W-1.5Zr), is used as a liner for high performance gun barrels because of its oxidation resistance and high-temperature strength. In addition, hafnium-based alloys with 20-27% tantalum and up to 2 % molyb-

8.4 Hafnium

denum are novel refractory alloys, exhibiting high oxidation resistance. The recent appearance on the market of a molybdenum-based alloy strengthened with hafnium carbide is another example of the pronounced ability of hafnium to improve creep resistance, hot strength, and recrystallization temperature. Intermetallics (see Chapter 11 in this Volume) are being studied as candidate materials for new-generation aircraft turbine engines. They are not only less dense than conventional superalloys, but some of them have the novel characteristic of exhibiting a yield strength which increases with increasing temperature. The Ni 3 Al compound is a promising intermetallic which can be solid-solution strengthened. However, only solutes like hafnium and zirconium are found to be effective for strengthening at temperatures above 800 °C. Although additions of about 2 %Hf slightly increase room-temperature yield strength, they increase the temperature at which peak strength is reached to about 750 °C. Resistance to creep, fatigue, and oxidation is also improved. Common iron-, nickel-, and cobalt-based precipitationhardened alloys also use some hafnium (Ayer et al., 1987, 1988). Since hafnium is a strong carbide former, hafnium-containing alloys become more resistant to stresscorrosion cracking. Under strong reducing conditions, carbon deposition and subsequent dissolution into heat-resistant alloys can occur at high temperatures. Therefore, internal carbide precipitates rich in chromium form in these alloys. Alloying with reactive elements such as hafnium is beneficial in reducing the carburization rates (Mitchell etal., 1992). A new, promising application for hafnium is in the construction of process equipment for spent fuel reprocessing

465

plants (Warnecke et al., 1976). Dissolution of spent fuel involves nitric acid solutions. The outstanding corrosion resistance of hafnium and its alloys towards this acid is complemented by a built-in safety factor due to hafnium's high neutron absorption capability. This unique combination of properties is applied to current designs for the chemical equipment in one reprocessing plant. In order to improve the design flexibility for neutron absorption and strength, TWC has studied the properties of Hf-Zr alloys and Hf-Ta alloys. Hf-Zr alloys (2.9%, 17.3%, 42.4%, 59.5%, and 81.4% Zr) and Hf-Ta alloys (1 %, 3%, and 5% Ta) have been evaluated for their corrosion resistance in various media. AH alloys have zero corrosion rate in boiling <70% HNO 3 with or without 1 % NaCl. Nevertheless, the Hf-59.5% Zr alloy seems to have better corrosion resistance than other Hf-Zr alloys in a solution of boiling 20% HCl + 200ppm Fe 3 + (TWC, 1987). Furthermore, the tensile strength of Hf-Ta alloys increases with increasing tantalum concentration. Average tensile strengths of Hf, Hf-1 % Ta, H f - 3 % Ta, and H f - 5 % Ta are 480, 520, 560, and 690 MPa, respectively. For a while, hafnium foil replaced zirconium foil as fuel in some photographic flashbulbs. Hafnium provides a higher intrinsic color temperature and a greater light output than zirconium. Furthermore, hafnium is being used in stereo headphones. It has also been used as a double carbide with niobium (60 Nb: 40 Hf), as a nitride coating for cutting tools, and as a minor addition to permanent magnet alloys. Pure hafnium is excellent for use as the active tip for plasma arc cutting tools. Recent technological advances in processing at TWC have made hafnium available at a higher purity level than was previ-

466

8 Titanium, Zirconium, and Hafnium

ously possible on a production basis. Specifically, the zirconium impurity level has been lowered routinely from 4.5% to less than 0.5 %. A zirconium impurity level of less than 0.1 % can be furnished when needed (Yau and Mark, 1989). Hence, new growing maskets for hafnium products have been opened up. For example, hafnium is replacing tungsten in coppercoated welding tips because of improved performance and durability. Other new applications include sputtering targets for several specialized electronic applications, and plasma arc cutting tips. Innovative areas of interest include the use of hafnium materials in vapor deposition and coating applications, such as coating of a carboncarbon composite with HfO2 for more durable oxidation protection. Hafnium oxide, a specialized refractory material, can be used to insulate thermocouples for short-term use above 2000 °C. Hafnium and hafnium oxide sputtering targets are used for coatings and specialized electronic applications. Hafnium tetrafluoride is used in some heavy-metal fluoride glass claddings.

8.4.7 Concluding Remarks and Future Thoughts

- At elevated temperatures, hafnium reacts with oxygen, carbon, and nitrogen to form stable, refractory hafnium oxide, carbide and nitride, respectively. These are unique compounds for specialized applications. - Hafnium is a very valuable alloying element. A small addition of hafnium can significantly enhance the behavior of many alloys, including superalloys, common stainless alloys, and permanent magnet alloys. - Hafnium can be alloyed with many elements to modify its mechanical, corrosion, and nuclear properties. As a result of these characteristics, hafnium has found applications in common objects, such as in stereo headphones, as well as in uncommon places, such as the space shuttle. So far, further development of applications has been restrained by a perceived limited availability and high cost. In fact, hafnium is no longer a scarce commodity. The availability and price stability of this metal and its compounds have spurred research and development efforts, leading to many new commercial applications. As the awareness of these materials and their properties becomes more widespread, an array of new products can be expected in the near future.

Hafnium is a unique metal. It possesses a set of novel characteristics: - It has a high capability to absorb thermal neutrons. This capability is about 600 times higher than in the case of zirconium. In this respect, hafnium and zirconium are complementary to each other in nuclear applications. - Hafnium is outstanding in its resistance to corrosion, oxidation, and irradiation damage. It resists attack by most acids, alkalies, saline solutions, molten salts, and gases.

8.5 References Abriata, J. P., Arras, D., Balrich, I C. (1983), Bull Alloy Phase Diagrams 4, 147. Arras, D., Abriata, J. P. (1988), Bull Alloy Phase Diagrams 9, 597. ASM (1990) Metals Handbook, Vol. 2,10th ed. Materials Park, OH: ASM Int., pp. 586-660. Ayer, R., Hayworth, H. C , Humphries, M. I (1987), European Patent 0 246 092. Ayer, R., Vaughn, G. A., Sykes, L. J. (1988), Exxon US Patent 4755240. Besson,- P., Kuhlman, P. U., Guerin, J., Brun, P., Bakes, M. (1976), US Patent 4021 531.

8.5 References

Bomberger, H. B., Froes, F. H., Morten, P. H. (1985), in: Titanium Technology: Present Status and Future Trends: Froes, F. H., Eylon D., Bomberger H. B. (Eds.), Dayton, OH: The Titanium Development Association, pp. 3-18. Boyer, R. R., Hall, J. A. (1993), Titanium '92, Science and Technology: Froes F. H., Caplan I. L. (Eds.). Warrendale, PA: TMS, Vol. 1, 77- 88. Christodoulou, L., Brupbacker, J. M. (1990), Mater. Edge 7, 29. Donachie, M. X, Jr. (1988), Titanium, A Technical Guide. Materials Park, OH: ASM Int. Driver, D. (1990), in: High Temperature Materials for Power Engineering: Bachelet, E., et al. (Eds.). Kluwer, Dordrecht, The Netherlands: Part II, pp. 883-903. Froes, F. H. (1991), Light Met. Age 49 (5/6), 6. Froes, F. H. (1994), J Mater. Set Technol. 10, 251. Froes, F. H., Baeslack, W. (1993), High Energy Electron Beam Welding and Materials Processing: Danko, J. C , Noltin, E. E. (Eds.). Miami Fl: American Welding Society, pp. 219-241. Froes, F. H., Eylon, D. (1990a), Int. Mater. Rev. 35, 162. Froes, F. H., Eylon, D. (1990b), Hydrogen Effects on Material Behavior: Moody, N. R., Thompson, A. W. (Eds.). Warrendale, PA: TMS, pp. 261-284 Froes, F. H., Hebeisen, J. (1994), Hot Isostatic Pressing '93: Delaey, L., Tas, H. (Eds.). Amsterdam: Elsevier, pp. 71-90. Froes, F. H., Suryanarayana, C. (1993), Reviews in Particulate Materials: Bose, A., German, R. M., Lawley, A. (Eds.). Princeton, NJ: Metal Powder Industries Federation, Vol. 1, pp. 223-275. Froes, F. H., Eylon, D., Bomberger, H. B. (Eds.) (1985), Titanium Technology: Present Status and Future Trends. Dayton, OH: The Titanium Development Association. Froes, F. H., Suryanarayana, C , Eliezer, D. (1991), ISIJ Int. 31, 1235. Froes, F. H., Suryanarayana, C , Eliezer, D. (1992), J. Mater. Sci. 27, 5113. Froes, F. H., Suryanarayana, C. Russell, K. C , Ward-Close, C. M. (1995), in: Proc. Int. Conf. on Novel Techniques in Synthesis and Processing of Advanced Materials, Rosemont, IL, Oct. 2-5, 1994: Singh, X, Copely, S. M. (Eds.). Warrendale, PA: TMS. Garzarolli, F , Stehle, H. (1987), in: Improvements in Water Reactor Fuel Technology and Utilization: IAEA STI/PUB/721, pp. 387- 407. Gee, E. A., Golden, L. B., Lusby, W. E., Jr. (1949) Ind. Eng. Chem. 41, 1668. Gegner, P. X, Wilson, W. L. (1950), Corrosion 15, 341t. Golden, L. B., Lane, I. R., Jr., Acherman, W. L. (1952), Ind. Eng. Chem. 44, 1930. Golden, L. B., Lane, I. R., Jr., Acherman, W. L. (1953), Ind. Eng. Chem. 45, 782.

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Guernsey, X B., Petersen, V. C , Froes, F. H. (1972), Met. Trans. 3, 339. Holzer, R., Stehle, H. (1986), CANDUNucl. Soc. Int. Conf Chalk River, Canada: AECL. Joseph, S. S., Froes, F. H. (1988), Light Met. Age 46, (11/12), 5. Kear, B. H. (1986), Sci. Am. 255(4), 158. Keller, H. W, Shallenberger, X M., Hollein, D. A., Hoth, A. C. (1982), Nucl. Technol, 59, 476. Kim, Y.-W (1994) JOM 46(7), 30. Lane, I. R., Jr., Golden, L. B., Acherman, W L. (1953), Ind. Eng. Chem, 45, 1067. MacKay, R., Bindley, P., Froes, F. H. (1992), JOM 43, (5), 23. Margolin, H., Neilson, X P. (1960), "Titanium Metallurgy", in: Modern Materials, Advances in Development and Applications, Vol. 2: Hauser, H. H. (Ed.). New York: Academic Press, pp. 225-325. Mitchell, D. R. G., Young, D. X, Kleeman, W. (1992), "The Effect of Reactive Element Additions on the Carburization Behavior of Heat-Resistant Fe-Based Steels", at First NACE Asian Conference, Singapore, Sept. 7-11. Houston, TX: NACE. Molchanova, E. K. (1965), Phase Diagrams of Titanium Alloys (Transl. of Atlas Diagram Sostoyaniya Titanovyk Splavov), Jerusalem: Israel Program for Scientific Translations. Morlin, K., Nordlof, S. (1985), "Performance Experience of ASEA-ATOM BWR Control Blades," in: Proc. Am. Nucl. Soc. Topical Mtg. on Light Water Reactor Fuel Performance, Orlando, FL, April 21 24. La Grange Park, IL: American Nuclear Society, Vol. 2, pp. 5.29-5.42. Moxson, V. (1993), Advanced Materials Products, Inc., Twinsburg, OH, private communication. Nielsen, R. H. (1980), "Hafnium and Hafnium Compounds", in: Kirk-Othmer: Encyclopedia ofChemcial Technology, Vol. 12, 3rd ed. New York: Wiley, pp. 67-80. Okamoto, H. (1992), / Phase Equilibria 13, 577. Polmear, I. X (1989), Light Alloys, Metallurgy of the Light Metals, 2nd edn. London: Edward Arnold. Schemel, X H. (1977), in: ASTM Manual on Zirconium and Hafnium, ASTM STP 639. Philadelphia, PA: ASTM, pp. 20-41. Schutz, R. W. (1994), JOM 46(3), 24. Seagle, S. R., Wood, X R. (1993), in: Synthesis, Processing and Modelling of Advanced Materials: Froes F. H., Khan, T. (Eds.): Aedermannsdorf, Switzerland: Trans. Tech, pp. 91-102. Smith, P. R., Froes, F. H. (1984), JOM 36(3), 19. Stickney, W. A. (1959), in: Zirconium-Hafnium Separation, United States Bureau of Mines, RI, p. 5499. Suryanarayana, C , Froes, F H., Rowe, R. G. (1991), Int. Mater. Rev. 36, 85. TWC (1987), Hafnium. Albany, OR: Teledyne Wah Chang. Upadhyaya, D., Ward-Close, C. M., Tsakiropoulos, P., Froes, F. H. (1994), JOM 46(11), 62.

468

8 Titanium, Zirconium, and Hafnium

Ward, C. (1993), Int. Mater. Rev. 38, 79. Ward-Close, C. M, Froes, R H. (1994), JOM46(1), 28. Warnecke, E., Comper, W, Poetzschke, M. (1976), "Use of Hafnium as a Heterogeneous Neutron Poison in Reprocessing Plants", at Reaktortagung, Dtsch. Atomforum e. V., Bonn, Germany, pp. 506509. Weidinger, H. G. (1983), in: IAEA Guidebook on Quality Control of Water Reactor Fuel, IAEA STI/ DOC/10/221. Vienna: AEA, pp. 111-131. Yau, T. L., Mark, W (1989), Outlook. Albany, OR: Teledyne Wah Chang, Vol. 10, No. 2, pp. 3-5. Yau, T. L., Webster, R. T. (1987), "Corrosion of Zirconium and Hafnium", in: Metals Handbook, Vol. 13: Corrosion, 9th ed. Materials Park, OH: ASM Int., pp. 707-721.

General Reading Titanium ASM (1990), "Properties and Selection: Nonferrous Alloys and Special Purpose Materials", in: Metals Handbook, Vol. 2, 10th ed. Materials Park, OH: ASM Int., pp. 586-660. Collings, E. W. (1984), The Physical Metallurgy of Titanium Alloys: Materials Park, OH: ASM Int. Donachie, M. X, Jr. (1988), Titanium, A Technical Guide. Materials Park, OH: ASM Int.

Froes, F. H., Caplan, I. L. (Eds.) (1993), Titanium '92, Science and Technology: Warrendale, PA: TMS. Froes, F. H., Eylon, D., Bomberger, H. B. (Eds.) (1985), Titanium Technology: Present Status and Future Trends. Dayton, OH: The Titanium Development Association. Kimura, H., Izumi, O. (Eds.) (1980), Titanium '80, Science and Technology. Warrendale, PA: TMS. Lacombe, P., Tricot, R., Beranger, G. (Eds.) (1989), Proc. 6th World Conf. on Titanium. Les Ulis, France: Les Editions de Physique. Lutjering, G., Zwicker, U., Bunk, W. (Eds.) (1985), Titanium '85, Science and Technology. Warrendale, PA: TMS. Zirconium and Hafnium ASM (1990), "Properties and Selection: Noferrous Alloys and Special Purpose Materials", in: Metals Handbook, Vol. 2, 10th ed. Materials Park, OH: ASM Int., pp. 661-669. Douglass, D. L. (1971), The Metallurgy of Zirconium. Vienna: Int. Atomic Energy Agency. Rickover, H., G., Geiger, L. D., Lustman, B. (1975), History of the Development of Zirconium Alloys for Use in Nuclear Reactors: U.S. Energy Research and Development Administration, Division of Naval Reactors, TID-26 740. Washington, DC: U.S. Gov. Printing Office. Schemel, I H. (1977), ASTM Manual on Zirconium and Hafnium, ASTM STP 639. Philadelphia, PA: ASTM. USBM (1956), Zirconium: Its Production and Properties, United States Bureau of Mines Bulletin 561.

9 Noble Metals and Their Alloys Giinther Schlamp

Formerly of Demetron (Degussa), Germany

List of 9.1 9.2 9.2.1 9.2.2 9.2.2.1 9.2.2.2 9.2.2.3 9.2.3 9.2.4 9.2.4.1 9.2.4.2 9.2.4.3 9.2.5 9.2.6 9.2.6.1 9.2.6.2 9.2.6.3 9.3 9.3.1 9.3.2 9.3.2.1 9.3.2.2 9.3.2.3 9.3.3 9.3.4 9.3.4.1 9.3.4.2 9.3.4.3 9.3.4.4 9.3.5 9.3.5.1 9.3.5.2 9.3.6 9.3.6.1

Symbols and Abbreviations Introduction The Elements History Occurrence Gold Silver Platinum and the Accompanying Platinum Group Metals Reserves and Resources Production Gold Silver Platinum Group Metals Commercial Grades and Purities Properties General Considerations The Atoms The Solid Metallic State Alloys General Remarks Phase Formation in Noble Metal Alloy Systems Primary Solid Solutions Intermetallic Compounds Superlattices Selected Phase Diagrams Influence of Alloying Elements Basic Considerations Metals Carbon in Platinum Group Metals Grain Refining Noble Metals and Gases Oxygen Hydrogen Liquid-State Properties General

Materials Science and Technology Copyright © WILEY-VCH Verlag GmbH & Co KGaA. Allrightsreserved.

471 474 475 475 476 476 476 476 477 478 478 479 480 481 483 483 484 486 497 497 497 497 500 504 507 510 510 510 512 513 514 514 516 517 517

470

9.3.6.2 9.3.6.3 9.3.6.4 9.3.6.5 9.3.7 9.3.7.1 9.3.7.2 9.3.7.3 9.3.7.4 9.3.7.5 9.3.7.6 9.3.7.7 9.3.7.8 9.4 9.4.1 9.4.2 9.4.2.1 9.4.2.2 9.4.2.3 9.4.2.4 9.4.3 9.4.3.1 9.4.3.2 9.4.4 9.4.4.1 9.4.4.2 9.4.4.3 9.4.5 9.4.5.1 9.4.5.2 9.5 9.5.1 9.5.2 9.5.2.1 9.5.2.2 9.5.2.3 9.6 9.6.1 9.6.2 9.6.3 9.7

9 Noble Metals and Their Alloys

Density Viscosity Surface Tension Electrical Resistance Special Alloys General Silver Alloys Gold Alloys Platinum Alloys Palladium Alloys Solders Dispersion-Hardened Alloys Composite Materials Manufacturing: Materials and Methods General Bulk Material Made by Melting Melting Conditions Casting Mechanical Processing Thermal Treatment Powders and Powder-Metallurgical Processing Powders Powder-Metallurgical Processing Textures General Classification Single Crystals Polycrystalline Materials Deposition of Coatings Coatings from Bulk Material Components Coatings from Chemical Compounds or Solutions Applications General Main Application Areas Non-industrial Applications Components and Connecting Materials Combined Physical and Chemical Interactions Recycling Direct Melting Dissolution Pyrometallurgical Processes and Equipment References

517 518 519 520 520 520 520 521 522 523 524 526 528 530 530 530 530 532 532 533 534 534 537 538 538 539 539 540 540 542 545 545 547 547 553 568 580 580 580 580 582

List of Symbols and Abbreviations

List of Symbols and Abbreviations A a, b, c C d e/a E G H Hc »o »r

K L M n N(r) P r ^opt

R S T Te Te Tm V

V X

a at A s s

n X

X V V

Q

Qo Qe

ampere lattice constants Curie constant grain diameter number of electrons/number of atoms thermoelectric voltage modulus of rigidity magnetic field strength coercivity intensity of incident light intensity of reflected light bulk modulus; constant length magnetic moment per mole number of electrons; refractive index occupation density pressure radius atomic radius optical reflectivity gas constant magnetostriction temperature transition temperature eutectic temperature melting temperature valency of solute valency of solvent; volume atomic percent solute temperature coefficient of electrical resistance coefficient of thermal expansion atomic diameter emissivity strain rate viscosity extinction coefficient thermal conductivity fraction optical frequency density average mass-density electrical resistivity

471

472

9 Noble Metals and Their Alloys

£ P h> Qi> Qd electrical resistivity due to phonons, ions, defects gr mass-density at radius r a surface tension a0 constant cre electrical conductivity (7 S yield strength T plateau stress TC critical resolved shear stress X susceptibility F o r symbols for crystal structures a n d basic types, see Volume 1 of this series: Structure of Solids. AAS AC AES APB b.c.c. B ET CIP CPC CVD DAB DC DNA e.m.f. ETAAS FAAS f.c.c. GGG HB h.c.p. HIP HT HV I ACS IC ICP-AFS ICP-MS m.p. LEED NM NTC OES PAFC PC

atomic absorption spectroscopy alternating current Auger electron spectrometry antiphase boundary body-centered cubic Brunauer - Emmett - Teller cold isostatic pressing controlled potential coulometry chemical vapor deposition Deutsches Arzneibuch (German purity standard) direct current deoxyribonucleic acid electromotive force electrothermal atomic absorption spectrometry flame atomic absorption spectroscopy face-centered cubic gadolinium gallium garnet Brinell hardness hexagonal close-packed hot isostatic pressing high temperature Vickers hardness International Annealed Copper Standard (see p. 291 of this Volume) integrated circuit inductively coupled plasma atomic fluorescence spectrometry inductively coupled plasma mass spectrometry melting point low energy electron diffraction noble metal negative temperature coefficient optical emission spectroscopy phosphoric acid fuel cell printed circuit

List of Symbols and Abbreviations

PCB PEM PES PFM PGM PMP PTFE PVD RE RRR SD SEM-EDX SIMS SPE SSMS TAB TCR UHF UPS VEC WRS XRF XPS YAG YIG ZGS

printed circuit board photon exchange membrane plasma emission spectrometry porcelain fused to metal platinum group metal powder-metallurgical processing poly(tetrafluoroethylene) physical vapor deposition rare earth residual resistance ratio sulfonylamido pyrimidine scanning electron microscopy energy dispersive X-ray secondary ion mass spectrometry solid polymer electrolyte spark source mass spectroscopy tape automated bonding temperature coefficient of electrical resistance ultrahigh frequency ultraviolet photoelectron spectroscopy valence electron concentration Wolffram's red salt X-ray fluorescence X-ray photoelectron spectrometry yttrium aluminum garnet yttrium iron garnet zirconium oxide grain-stabilized

473

474

9 Noble Metals and Their Alloys

In the periodic table of the elements they are the middle and bottom members of the transition metals in groups 8 and 1B of the 5th and 6th period (see Fig. 9-1). The noble metals represent typical metallic elements with very dense, close-packed crystal structures, high densities, brightness, high electron mobilities, ductility, and the capability to form alloys. Based on their individual electronic structures, they show, unique to each, characteristic physical and chemical properties, covering a wide range in respect of melting temperature, hardness, or electrical conductivity. The elements with the highest density (iridium 22.61 g/cm3), the highest electrical conductivity (silver 6.14 Q,"x cm" 1 ), and the highest resistance to oxidation (gold: standard electrode potential: + 1.5V) are in the noble metals group. Systematic relations exist for the shift of properties along the periods as well as inside each group, including the corresponding elements of the 4th period. Examples include the alloying behavior of silver and gold compared to copper, the formation of magnetic alloys, the catalytic activities of

9.1 Introduction Noble metals comprise the elements gold, silver, and the six platinum group metals (PGM) ruthenium, rhodium, palladium, osmium, iridium, and platinum. Gold and silver were appraised in ancient times as precious, high value metals, primarily used for jewelry, coins, silverware, and in dentistry. In addition to these still continuing applications, noble metals and their alloys have become important key materials in numerous different technical fields especially after the discovery of platinum and the separation and growing availability of the platinum group metals in the 19th century. They are of equally high technical and commercial importance as constituents and components of devices and production materials for electrical engineering, electronics, and sensor technology, as catalysts for chemical processing and as automotive catalysts, as electrodes of galvanic, battery and fuel cells, as heatreflecting coatings on architectural glasses, and as highly coercive magnetic alloys for recording and memory disks.

1 1A

2A 4

3

Li 11

Na K 37

Rb 55

Cs

3B

HP

B

Be

13

12

Al

Mq 20

Ca 38

Sr 56

4A 22

21

Sc

Ti 40

39

Y

Zr 72

S7-71

Hf

Ba

5A 23

V 41

Nb 73

Ta

6A 24

7A 25

43

Mo 74

26

Mn Fe

Cr 42

8

Tc 75

W

Re

44

Ru 76

Os

27

Co 45

Rh 77

Ir

1B ~28

Ni 46

Pd 78

pt

1

29

Cu 47

Ag 79

Au

87

Fr

88

Ra

89-'J03 \104

1

/ Ac tin ides

105

5B 7

C 14

Si

6B 8

M

N 15

P

0 16

s

7B 9 r

r 17

10

Ne 18

Ct

Ar

2B 30

Zn 48

Cd 80

Hg

31

Ga 49

In 81

Tl

\ 1

4B 6

5

3A 19

2

H

106

Ha \

Lanthanides

Figure 9-1. Position of the noble metals in the periodic table.

32

Ge 50

Sn 82

Pb

33

As 51

Sb 83

Bi

34

Se 52

35

Br

84

Po

Kr 54

53

Te

36

I 85

At

Xe 86

Rn

9.2 The Elements

the platinum group metals compared to those of iron, nickel, and cobalt, and the formation of complex chemical platinum group metal compounds, which have become especially important in the synthesis of stereo-enantiomer pharmaceuticals. In both heterogeneous and homogeneous catalysis, the platinum group metals and alloys perform more efficiently and with higher selectivity than those of the corresponding 4th group elements.

9-2 The Elements 9.2.1 History Gold's existence was known even in prehistoric times. The name derives from the Indo-German word meaning yellow and bright. Found in elementary form, it was collected and worked manually to cult, ornamental and jewelry objects, and as bars and coins as the first kind of money. Manufacturing centers arose between 6000 and 3800 B.C. in Egypt, Greece, and the Near East, along similar lines to copper and silver metallurgy. New important deposits were found in the 1700s in South, Central, and North America, Alaska, Russia, Australia, South Africa, New Zealand, and New Guinea. Silver's existence has also been known since about 6000 years B.C. and has been of great cultural importance ever since. The origin of its name is unknown; the name silver is derived from a Teutonic source; the symbol Ag, from the Latin word argentum. It was used like gold for decorative and jewelry objects, and later in large quantities for silver coins (Egypt, Greece, China, and Japan). Large silver deposits in Europe were exploited in the middle ages. Today's production centers are located in South America, Mexico, Canada, and the U.S.A.

475

The first discovery of platinum can be traced to the Mayas in Central America in about 1550 A.D., when European colonists entered the continent. Looking for gold and silver, they disregarded the hard gray metal as unimportant and unwanted and named it "platina" (Spanish meaning little silver). In 1751 it was recognized in Europe as a true metal (Watson, 1751). Its first practical use was in jewelry as a harder precious metal. Platinum gained growing importance from the mid 1800s, when it became available in industrial quantities. The development of technical production processes (sintering, melting and mechanical forming) was stipulated due to increasing interest from electrical, dentistry, and chemical technology fields. The development of chemical analysis in the 1800s led to the discovery, separation, and production of the platinum-accompanying noble metals palladium, rhodium, iridium, ruthenium, and osmium. The two most important - palladium and rhodium - only became available in technical quantities in 1930, when notable platinum production started with the exploitation of large nickel ore deposits in South Africa and Canada. • Palladium was discovered in 1804 (Wollaston, 1804) and named after the newly discovered, asteroid "Pallas", meaning safeguard of liberty. • Rhodium was also discovered by Wollaston (1805) and named after the Greek word "rhodon" because of the rose-red color of its compounds. • Iridium was discoverd by Tennant (1804) and named after the Greek goddess of the rainbow, Iris, because of the brillant color-effect of its salts when dissolved in HC1. • Osmium was discovered by Tennant (1804), and named after the Greek word

476

9 Noble Metals and Their Alloys

"odor", as its compounds have a pungent, disagreeable smell. • Ruthenium was observed in 1828 in platinum ores in the Ural mountains, and was first recognized as an element by Claus (1845). Its name is the old Latin name for Russia (Ukraine). 9.2.2 Occurrence 9.2.2.1 Gold Gold is distributed very unevenly in the earth's crust due to enriching processes that have taken place near the surface. The average abundance is estimated as ^0.005 ppm (parts per million). The gold content in the oceans is assumed to be ^0.008-4 mg/m 3 (i.e., ppb, parts per billion). The sources contain gold as the metal alloyed with varying amounts of silver or copper. Primary deposits consist of quartz veins or pyrites with incorporated gold particles, formed by hydrothermal processes with copper ores. Secondary deposits (placers) are enriched deposits formed by the wear of rocks by weathering and consolidation due to flowing water (river sands). The richest and most important deposits at present are conglomerates with sand and shingle, which contain up to 12 ppm gold when separated from the gangue. The most important deposits are in South Africa (Witwatersrand), Russia, the U.S.A., Papua New Guinea (Ok Tedi), and Brazil. World gold reserves are assessed at 70 000 t. 9.2.2.2 Silver In the region accessible by mining (to a depth of 5 km), the silver amounts to an average of ^0.5-0.1 ppm, presently at least 100 times higher than that of gold or

platinum. Also, in the oceans the abundance of silver is estimated to be much higher, ^0.001 ppm (^10 9 t), which cannot however be extracted economically. Silver became concentrated at particular sites in the earth's crust by the recrystallization of silicate magmas caused by volcanic action in impregnation zones. Native silver occurs in the form of nuggets, dendrites, or lumps. Main sources are silver sulfides in complex ores with lead, leadzinc, to a lesser amount copper and nickel, selenides, tellurides, arsenides, antimonides, and bismuth compounds. Argentiferous copper ores satisfy ^ 3 0 % of the world silver demand. Even very low contents of silver are profitably recovered, because during the electrorefining of copper the silver is collected in the anode slimes. The extraction of silver from gold ores results in less than ^ 1 0 % , and extraction from tin deposits ^ 2 % or world production. Argentiferous lead and lead-zinc ores are located in North America, Mexico, South Africa, Russia, and Australia. Silver-containing copper ores are found in Europe, Canada, Chile, and South Africa. The world silver reserves amount to ^250 000 t. 9.2.2.3 Platinum and the Accompanying Platinum Group Metals The abundance of the PGMs is expected to be of the order of 10 ~ 4 ppm. They are concentrated in planetary regions reaching up to «30 ppm in the earth. Due to the mainly siderophilic chemical character of the PGMs, virtually their entire mass seems to be in the earth's metallic core. The siliceous lithosphere is estimated to contain 0.05-0.5 ppm. As with gold, there are primary and secondary deposits. Primary deposits are found as metallic (alloyed) and sulfidic in-

9.2 The Elements

elusions in carrier magnesium-iron silicate minerals. Secondary deposits originate from oxidation processes of PGM-bearing minerals or the erosive action of water forming river headwater placers. Typical nuggets or lumps were formed by mechanical agglomeration involving chloridic dissolution and reprecipitation. The composition of minerals in regard to PGM-mix vary according to their provenance (Table 9-1). Special nickel ores contain higher amounts of palladium, whereas osmium and iridium occur at higher concentrations in minerals like osmiridium and iridiumplatinum. The largest deposits of PGMs are found in South Africa (Bushveld: Merensky and, UG 2 Reef), in the U.S.A. (Stillwater), in Canada (Ontario), and in Russia. The latter two are the most important sources for palladium. The world reserves of PGMs are estimated to be «70 000 t. The distribution of relative amounts of the PGMs in the earth is estimated to be 20% each of platinum, palladium, ruthenium, and osmium, and 6% each of rhodium and iridium. These numbers differ considerably from the typical compositions of the ores found in different areas.

477

9.2.3 Reserves and Resources

Estimates on the reserves (which affirm those resources whose existence has been estabilished by projecting and for which mining is economically viable) and on additional resources (for those with no current economic value) depend on technical and economical parameters. Today, gold reserves, which are economically viable, are assessed at 60000 tons (60 000 tones), which represent more than 50 times the world production per annum from primary deposits. Twenty years ago the gold reserves were calculated as only 20% of this amount. At that time the extensive deposits in Brazil had not been discovered, and the profit margin was low due to the fluctuating gold price. Of the present gold reserves 90% are found in three countries: in the Republic of South Africa (40%), in Brazil (35%), and in the Commonwealth of Independent States, C.I.S. (15%), followed by the U.S.A., Canada, Australia, Venezuela, Zimbabwe, and Ghana with amounts from 1 to 3%. The corresponding silver reserves are distributed more evenly over the entire globe. At the present time the mineral re-

Table 9-1. Composition of PGM minerals (Renner, 1992). Bushveld complex -

Platinum (%) Palladium (%) Ruthenium (%) Rhodium (%) Iridium (%) Osmium (%) Gold (%) Grade (g/t)

Merensky Reef

UG2 Reef

59 25 8 3 1 0.8 3.2

42 35 12 8 2.3

8.1

0.7 8.71

South Africa, Plat Reef

Sudbury, Canada

Noril'sk, CIS

Colombia

Stillwater, U.S.A.

Average

42 46 4 3 0.8 0.6 3.4

38 40 2.9 3.3 1.2 1.2 13.5

25 71 1 3

93 1

19 66.5 4.0 7.6 2.4

45 30 5 4 1 <1

7-27

0.9

2 3 1

0.5 3.8

22.3

478

9 Noble Metals and Their Alloys

serves amount to 250 000 t of silver, which represents about 20 times the annual world production. By far the largest proportion of silver is found in the C.I.S. (^25%), in the U.S.A. 0 2 0 % ) , in Mexico («15%), and in Canada and Peru (^10% each), followed by Poland, Chile, Brazil, Australia, Zaire, the Republic of South Africa, Spain, Sweden, former Yugoslavia, Japan, and China with minor quantities. The largest known primary deposit for the platinum group metals is in the Republic of South Africa. The reserves amount to 70% for combined PGMs and 95% for rhodium. The Bushveld Complex in South Africa is estimated to be inexhaustible. Next in line are the U.S.A., the C.I.S., and Canada with approximately 12 to 5%. The Stillwater complex in Montana and the nickel deposits in Siberia and in Ontario are important sources in particular for palladium. Smaller reserves are reported for Australia, Columbia, and Zimbabwe. The current estimate of workable deposits with accessible economic value is 70 0001, which represents a 300-fold tonnage of the current world production.

means such as milling, flotation, or panning. The most important process for gold extraction is cyanide leaching in the presence of particles of activated carbon ("carbonin-pulp" process). The gold cyanide solution is adsorbed on the surface of the carbon particles. To minimize losses the process of solution and adsorption runs in several stages in countercurrent flow (Fig. 9-2). The loaded carbon particles are removed from the pulp by sieving; the adsorbed gold cyanide can be eluted by washing with hot NaCN solution. Gold is recovered by either precipitation with zinc or electrolysis with stainless steel electrodes. The carbon regeneration is carried out by washing with acid to remove calcium carbonate and base metals, and by thermal reactivation to remove organic materials and reexpose the pore structure. New developments are directed to exchanging the carbon for organic resins as the adsorbing agent agent for the AuCN solution ("resin-in-pulp"). Depending on the grade and purity, two different processes are used:

9.2.4.1 Gold

1) The Miller process uses chlorine gas passed over the surface of molten gold at ^1100°C. At this temperature, silver and base metals form stable chlorides. AgCl and CuCl can be removed as molten slag, the chlorides of the other metals evaporate leaving gold of purity 99.5-99.6%. 2) Wohlwill electrolysis used an electrolyte containing HC1 and H(AuCl4) in aqueous solution. The cathode consists of titanium sheets, the anodes are made from raw gold, leading to fine gold of 99.999.99% purity. The preferred anode material is gold which has already been refined by the Miller process.

Gold particles are concentrated and separated from the gangue by mechanical

Solvent extraction is based on the different solubilities of noble metal salts in

9.2.4 Production Processes for mining, production, and refining of the noble metals have been developed in close correlation to the progress of industrial technology and demand and requirements for applications. Production processes of noble metals include concentration (separation from the gangue), extraction, recovery, and refining. Detailed process steps depend on the type of mineral, particle size, and associate materials.

9.2 The Elements

479

Continuous process (computer-controlled)

k

Au cyanide solution + ore pulp Carbon + pulp

jO.O1gAu/t)

1

Carbon fines* for recycling

Carbon+ pulp Carbon loaded with Au

(0.7gAu/t)

Pulp recycle Carbon (reactivated) for recycling

Na(Au(Cn) 2 sol.

Carbon (Au-free)

Dump Raw Au for refining

NaCN solution (0.02gAu/t)

Figure 9-2. The carbon-in-pulp process: 1) mixing tanks, 2) screens, 3) elution (with NaCN solution at 110°C), 4) reduction electrolysis, and 5) reactivation (at 750 °C), (Renner).

aqueous and in organic solvents. By forming complex or ion-pair compounds, noble metals are enriched in the organic phase, which is not miscible with the aqueous solution and can be extracted by separating the liquid phases. Solvent extraction offers the possibility of separation with quantitative yield and highly selective action. Gold can be effectively separated in the solvent extraction process with dibutylcarbitol (Fiberg and Edwards, 1978). 9.2.4.2 Silver True silver ores are treated by the aqueous cyanide process, which is similar to that use for gold leaching. The ore is crushed, wet milled, and leached with an aqueous solution of NaCN. Solids are removed by filtering. Silver is then precipitated either by solidification with zinc dust or by electrolytic reduction. Adsorption of the cyanide complex on activated carbon is also possible but of lesser importance. By far most of the silver comes from lead and zinc ores, from which it is concen-

trated together with gold and the platinum metals by the Parkes process. By stirring 1.5% zinc in the silver-containing lead melt, Ag-Zn crystals are formed which separate from the melt; this crust contains «10% silver. The zinc is distilled leaving a 50% silver concentrate with predominantly lead. From this, silver is extracted by the Cupellation process. In this process, air is blown onto the liquid melt at temperatures above the melting temperature of PbO. The formed liquid PbO layer dissolves oxides of the base metals and up to 5% silver. It is continuously withdrawn for renewed extraction, leaving highly concentrated silver of ^ 9 9 % purity. Silver from silver-containing copper ores is collected in the anode slimes of electrolytic copper refining. It is extracted by either hydrometallurgical or pyrometallurgical processes (see Sect. 9-6). Refining is done by electrolysis using the Moebius process. It works with a silver nitrate-sodium nitrate-nitric acid elec-

480

9 Noble Metals and Their Alloys

trolyte, stainless steel cathodes, and cast, crude silver anodes. Silver is deposited on the cathodes and continuously removed. The remaining slime contains gold and PGMs. A purity of up to 99.99% is obtained with a single run, 99.999% with a double run. 9.2.4.3 Platinum Group Metals Production steps for the PGMs include sequences of mechanical, pyrometallurgical, and hydrometallurgical processes (Fig. 9-3). After flotation, the material is melted under alternating reducing and oxidizing conditions, removing iron and SO 2 , followed by grinding and then magnetic separation of the enriched PGM-containing Ni-Cu-Fe phase. Solution with aqua regia or HC1/C12 treatment removes most of the base metals, leaving a PGM concentrate of above 30% PGM. Osmium and ruthenium are first separated by distillation of their volatile oxides from appropriate concentrated hydrochloric solutions. The remaining PGMs are then separated in aqueous solution by subsequent oxidation/reduction and precipitation of complex PGM chlorides in the sequence: platinum, iridium, palladium, and rhodium (Fig. 9-4). Instead of precipitation from aqueous solutions, PGMs can also be recovered from solutions by solvent extraction. This gives a quantitative yield and is highly selective. Refining is done by repeated solution, recrystallization, and precipitation of compounds. Conversion to metal, sponge, or powder is carried out for platinum, palladium, iridium, and rhodium by thermal decomposition of the ammonium chlorocomplexes (calcination) in an inert gas atmosphere or a hydrogen-nitrogen mixture.

Ore 5 - 2 0 g / t PGM

Grinding Flotation

Flotation concentrate { 1 6 0 - 2 0 0 0 g / t PGM)

n

Reductive smelting of an .Ni-Cu-Fe matte

-Slags

Green matte -Slags-

Oxidation of matte Separation of iron Adjustment of sulfur content

-S0,

White matte (low iron)

Slow cooling -Low-PGM Cu-Ni matte

Grinding Magnetic separation Ni-Cu-Fe phase

| Dissolution [

*• Ni-Cu-Fe solution

PGM-concentrate ( > 3 0 % PGM)

| Removal of base metals |

| PGM separation |

Platinum group metals

Figure 9-3. Production steps of PGMs (Ullmann's, 1992).

Ruthenium is prepared by precipitating RuO 2 from aqueous solution and reducing it using hydrogen, osmium by precipitating potassium osmate(iv) and heating it in a hydrogen atmosphere or by reduction with NaBH 4 in an aqueous solution.

481

9.2 The Elements PGM solution, oxidized

(prci 6 r,(Pdci 6 r, (Rhci6r, drci 6 r

I

Boiling, Pd reduction ( P d C l 6 ) 2 " - —•»•- ( P d C l 4 ) 2 " - • C l 2

Ir reduction (IrCl6)

2-

(IrCl6)3"

Pt precipitation 2

(PfCl6) "+ 2NH4+

(NH4)2(PtCl6)

Oxidation ( P d C l 6 ) \ ( I r C l 6 r , (RhC(6) Boiling, Pd reduction (PdCL

(PdCU

Cl2

Ir precipitation 2

(IrCl6) ~+ 2NH4+

(NH4)2(IrCl6)

Oxidation (PdCl4)2~+Cl2

(PdCl 6 )'

Pd precipidation 2

(PdCL) ~+2NH,+

(NH 4 ) 2 (PdCl 6 )

use. One carat corresponds to 1/24 of the total metal, a total of 24 carat representing 100% pure noble metal. Twenty-four carat equals 1000 fineness, 18 carat corresponds to 750 fineness (Table 9-2). The noble metal contents of ores and intermediate products are given in grams per metric tonne (i.e., ppm) or ounces per metric tonne, or corresponding units. The minimum content of noble metal and hallmarking of defined fineness grades are regulated by law in most countries, but differ to a considerable extent (e.g., for gold alloys from 800/1000 in Portugal, through 750/1000 in France, 585/1000 in most European countries, 416/ 1000 in the U.S.A., and 333/1000 in Denmark, Germany, Italy, and Greece). The International Organisation for Standardisation has recommended [ISO 9202, 1991 (E)] standardized values to be applied in the European Community (E.C.) (Table 9-3). Standard commercial grades together with the maximum permitted content of each impurity element have been designed by the ASTM (Table 9-4). Typical analyses

Concentration

A.

Rh precipitation (RhCL

(NH 4 ) 3 (RHCl 6 )

Figure 9-4. Separation of PGMs (Gerner et al, 1991).

9.2.5 Commercial Grades and Purities

Noble metals are generally applied in high purity grades as pure elements and as alloys. Commercial grades start with the "good delivery" quality for gold (99.5%) and silver (99.9%), and "technical pure" for platinum metals (99%). The silver and gold content of pure or alloyed metal in bulk or surface layer form is expressed in fineness (parts per thousand). For gold the older designation of carat is also still in

Table 9-2. Fineness and corresponding carat ofjewelry alloys. Fineness in 1/1000

Corresponding carat

333 375 585 750 1000

3 9 14 18 24

Table 9-3. Standard fineness of noble metal alloys3. Metal

Fineness (Standard compositions)

Au Ag Pt Pd

375

585

500 500

850

750 800 900

According to ISO 9202 (1991) E.

916 925 950 950

999 999 999 999

482

9 Noble Metals and Their Alloys

Table 9-4. ASTM purity grades. ASTM designation

Grade

Ag

B 413-69 (reappraised 1981)

99.9 99.95 99.99 99.999

Au

B 562-86

99.5 99.95 99.99 99.995

Pd

B 589-82 (reappraised 1987)

99.8 99.95

Pt

B 561-86

99.8 99.95 99.99

Rh

B 616-78 (reappraised 1983)

99.8 99.9 99.95

Ir

B 671-81 (reappraised 1987)

99.8 99.9

Ru

B 717-

99.8 99.9

Metal

of silver and gold grades are given in Table 9-5. For most technical applications, silver is required with a fineness of 99.97% (e.g., for alloying) or 99.99% (for electrical contacts and electroplating). Highest purities of up to 99.999% are needed for gold, silver, and platinum in electronic applications. Analytical control of the precious metal content in deposits, ores, and recycling material, and purity analysis of the refined grades are of high importance both in view of their monetary value and regarding purity requirements for technical applications. Special analytical procedures, based on various metallurgical, chemical, and physical processes, have been developed. The most common metallurgical method for the determination of gold and silver is the so-called fire assay, also called "docimasy". The method consists essen-

tially of the extraction of the noble metals by melting the samples together with lead at temperatures between 800 and 1200°C under oxidizing conditions. Gold and silver are dissolved and separated, while the base metals are converted to slag. Chemical methods start with solution of the samples, followed by precipitation of the precious metals by element specific reagents, and gravimetric evaluation. Numerous modern physical methods with highly sophisticated instrumentation are especially suited for trace analysis of high grade materials. Starting with atomic absorption spectroscopy (AAS), highly specialized spectroscopic methods include optical emission spectroscopy (OES), spark source mass spectroscopy (SSMS), X-ray fluorescence (XRF), flame atomic absorption spectroscopy (FAAS), electrothermal atomic absorption spectrometry (ETAAS), plasma emission spectrometry (PES), inductively coupled plasma mass spectrometry (ICP-MS), and inductively coupled plasma atomic fluorescence spectrometry (ICP-AFS). Further methods include activation analysis, controlled-potential coulometry (CPC), polarographic differential pulse polarography, anodic stripping voltammetry, ion-selective electrode potentiometry, and ion chromatography. A special field is the examination of the surface conditions needed for research on catalytic activities. Suitable methods are X-ray photoelectron spectrometry (XPS), scanning electron microscopy energy dispersive X-ray detection (SEM-EDX), Auger electron spectrometry (AES), secondary ion mass spectrometry (SIMS), ultraviolet photoelectron spectroscopy (UPS), and low energy electron diffraction (LEED). The quantitative determination of trace elements requires the preparation of calibrated standards (Furuya, Kallmann 1991).

483

9.2 The Elements

Table 9-5. Commercial (a) gold and (b) silver grades (a) (Renner, 1990, 1)

(b) (Renner, 1993, 1)

Designation

Designation

Good delivery gold Fine gold

Fine gold, chemically pure Fine gold, high purity

Gold content (w.%)

Content of other metals (ppm)

>99.5

any metals (total) <5000 <100 > 99.99 Ag Cu < 20 others < 30 total <100 < 25 > 99.995 Ag others < 25 total < 50 > 99.999 Ag < 3 Fe < 3 Bi < 2 Al < 0.5 Cu < 0.5 Ni < 0.5 Pd + Pt < 5.0 total < 10

9.2.6 Properties

Impurities (ppm)

Good delivery silver Fine silver

Fine silver 999.9

Fine silver, chemically pure

Fine silver, high purity

9.2.6.1 General Considerations

The electronic structure and size of the atoms determine the type and strength of the interatomic bonding forces, the melting temperatures, and the chemical behavior (oxidation states, oxidation resistance, stability of alloys and compounds). The electrical and thermal conductivity, and the magnetic and optical properties are essentially determined by the individual electronic band structures of the elements in the solid metallic state, the free, movable electrons, and their interactions with the lattice and with phonons. The physical and chemical properties of the noble metals derive from their high nuclear charges and their complex electronic configurations. The noble metal are primarily characterized by their high resistance to oxidation and their high thermal and mechanical sta-

Silver content

Silver nitrate DAB 7, for photographic usea

a

<900 copper <100 others <300 > 99.97 copper < 50 others including < 10 lead bismuth < 10 <300 total < 50 < 99.99 copper < 50 others including 10 lead 10 bismuth 15 < 99.995 copper 35 others including 5 lead 5 iron 1 gold 1 bismuth 2 < 99.999 aluminum 2 iron 2 silicon 2 tin 1 lead 1 gold 0.5 cadmium 0.5 bismuth 0.4 copper 0.4 manganese 10 total 2 63.5 copper 2 iron 10 lead 10 sulfate 10 nitrate water 10 insoluble not precipitated byHCl < 20 moisture <200 >99.9

DAB = Deutsches Arzneibuch (German purity standard).

484

9 Noble Metals and Their Alloys

bility. Characteristic properties are given in Table 9-6. 9.2.6.2 The Atoms Electronic Structure The electronic structure of the transition metals is characterized by the complication of their 3d-, 4d-, and 5d-electron levels with increasing nuclear charge along the periods. The occupation of outer d- and s-shells reaches, for the "light" noble metals (at. no. 44-47), based on a completely filled [Ar]-shell, from 4d 7 5s 1 for ruthenium to 4d l o 5s 1 for silver, and for the "heavy" noble metals (at. no. 76-79), based on a completely filled [Kr]-shell, from 5d 6 s 2 for osmium to 5d lo 6s 1 for gold. Because of the very small energy differences between their outer s- and d-electron levels, d-electrons can easily be transferred into the valence band, therefore enabling noble metals to react in many different oxidation states. They form numerous solid, metallic solutions, alloys, and intermetallic and complex inorganic and organic compounds. The "noble character" is defined by the resistence to oxidation, which is measured by the reduction potential, and by the chemical stability, which is measured by the first ionization energy. The reduction potential and first ionization energy increase with increasing nuclear charge along periods and inside groups. They are determined by the degree of completion of the d-levels up to the stable d 10 configuration, and by the attraction of the outer s-electrons by the nuclear charge. The tendency to complete the occupation of the 4d- and 5d-levels can also be seen by the fact that - with the exception of osmium and iridium - the outer s-levels are only singly occupied. Palladium with its basic configuration of d l o s° thereby surpasses

silver in its reduction potential. The increase of the noble character with growing atomic number is essentially caused by a decrease of screening of the nuclear charge by the inner electrons. This enhances the inner-atomic bond strength, extending onto the valence electrons. Differences in the screening of nuclear attraction also explain the similarity in the chemical behavior of copper and silver compared to gold, and ruthenium and osmium, rhodium and iridium, and palladium and platinum compared to the 4th period elements iron, cobalt, and nickel. The screening of nuclear attraction is also a structure stabilizing factor of many complex chemical noble metal compounds. Atomic diameters Atomic sizes of the transition metals decrease with increasing atomic number in each period, beginning at group 3, and show minimum values in each period in group 8A (Fe, Ru, Os). The atomic diameters increase gradually on passing from ruthenium to silver and from osmium to gold, and then with a marked increase from palladium to silver and from platinum to gold. The atomic diameters of the transition metals of the second long period are systematically larger than those of the analogs of the first period. Caused by completion of the 4f-electron levels of the rare earth elements ("lanthanides contraction"), the atomic diameters of the elements of the third long period are very close to those of the group analogs of the second long period. The pairs ruthenium/ osmium, rhodium/iridium, palladium/platinum, and silver/gold have practically the same atomic diameters.

9.2 The Elements

485

Table 9-6. Properties of noble metals. Property

Noble metal Ru

Rh

Pd

Ag

Os

Ir

Atomic number Atomic weight Electron configuration

44 101.07 [Kr]4d75s

45 102.905 [Kr]4d85s

46 106.42 [Kr]4d10

47 107.868 [Kr]4dlo5s

Crystal structure Lattice constant (A) (lO^nm)

h.c.p. a =2.701 c =4.275 c/a = 1.583 1.325

f.c.c. 3.803

f.c.c. 3.887

f.c.c. 4.086

77 192.22 [Xe]4f14 5d76s f.c.c. 3.839

78 195.08 [Xe]4f14 5d96s f.c.c. 3.924

79 196.966 [Xe]4f14 5dlo6s f.c.c. 4.078

1.345

1.376

1.445

76 190.2 [Xe]4f14 5d66s2 h.c.p. a =2.735 c =4.319 c/a = 1.579 1.377

1.357

1.377

1.442

12.41

12.02

10.49

22.61

22.65

21.45

19.32

10.70 130 386 153 0.26 0.36

10.49 50 124 51 0.39 0.52

9.35 26 82 27 0.37 0.95

20.1 300-680a 570 220 0.25 0.28

19.39 200 538 214 0.26 0.28

18.91 48 173 67 0.39 0.36

17.19 25 79 28 0.42 0.58

2310

1966

1554

961.9

3045

2410

1772

1064.4

5.5 3900

3.5 3700

20 2970

1.5 2212

20 5000

4 4130

0.1 3827

0.02 3080

117

150

75

418

87

147

73

310

9.5

8.5

11.1

19.0

6.5

6.8

9.0

14.3

6.7

4.34

9.725

1.465

8.2

4.7

9.825

2.0

4.0

4.6

3.8

4.1

4.2

4.3

3.92

4.03

0.47

0.9

0.71

0.11

0.684 4.52

0.70 4.8

-0.570 4.97

0.740 4.73

4.55

0.660 5.40

0 5.27

0.770 5.22

8, 6, 4, 3, 2, 0, - 2 7.37

5, 4, 3, 1, 2,0 7.46

4, 2, 0

2, 1

3, 1

7.576

6 4, 3, 2, 1,, 0 , - 1 9.1

4, 2,0

8.34

8, 6, 4, 3, 2, 0, - 2 8.7

9.0

9.225

125

125

128

134

126

127

130

134

Pt

Au

Atomic properties

Atomic radius (A) (10" 1 nm) Mechanical properties

12.45 Density (20 °C) (g/cm3) Density, liquid at 3 melting point (g/cm) 10.90 Hardness (HV, kg/mm2) 250-5003 485 Young's modulus (GPa) Modulus of rigidity (MPa) 172 Poisson's ratio 0.29 0.31 Coefficient of compressibility Thermal properties

Melting point (m.p.) (°C) Vapor pressure at m.p. (mbar) (lxlO 5 Nm~ 2 ) Boiling point (°C) Thermal conductivity (20°C)(Wm- 1 K- 1 ) Coeff. of linear expansion (0-100°C)(10" 6 K) Electrical properties

Specific electrical resistance (0°C)(nQcm) Temperature coefficient of electrical resistance (0-100°C) (10- 3 K) Transition temperature of superconductivity Thermoelectric voltage against Pt (0-100 °C) (mV) Work function (eV) Chemical properties

Oxidation states in compounds Ionization energy (eV) Covalent radius for single bonds (Pauling) (pm) Ionic radius (pm) (oxidation number, coordination number) Electronegativity (Allred and Rochow) Reduction potential (V) with (n) electrons a

62(4 + , 6) 55 (5 + ..6) 64(2 + , 4) 79(2 + , 4) 39(8 + , 4) 63 (4 + ,6) 63 (4 + ,6) 68(3 + , 4) 68(3 + , 6) 67(3 + ,,6) 86 (2+ ,6) 67(1 + ,2) 63 (4 + ,6) 68 (3 + ,6) 60(2 + , 4) 137(1+,6) 1.4

1.5

1.4

1.4

1.5

1.6

1.4

1.4

0.455 (2)

0.758 (3)

0.951 (2)

0.800 (1)

0.85 (8)

1.156 (3)

1.118 (2)

1.498 (3)

Depending on crystal orientation.

486

9 Noble Metals and Their Alloys

9.2.6.3 The Solid Metallic State Electronic Structure In the condensed metallic phase, the discrete electron energy states of the outer s- and d-levels are broadened to energy bands, forming free, movable electrons in the periodic lattice of positive ions. Bragg reflection of the electrons on the lattice planes caused partial extinction, which leads to the formation of separate, occupied zones. These are represented in wavenumber space ("&-space") by the Brillouin zones, which derive directly from the crystal structure. The energy states occupied by the free electrons in the ground state fill up the Fermi sphere, which represents the maximum possible energy (the Fermi energy) to be taken by an electron inside the Brillouin zone at 0 K. The surface of equal energy in three-dimensional A>space is the Fermi surface. The electrical, magnetic, and optical properties of metals depend to a important degree on the behavior of the electrons in the region of the Fermi surface and electron transitions between occupied states below and unoccupied states above the Fermi energy (Kittel, 1986). The Fermi surfaces of silver and gold are nearly spherical except for the ringshaped intersections centered on the faces of the octahedron, forming necks, caused by the 5s/4d or 6s/5d interactions (Fig. 9-5; Christensen, 1972; Christensen and Sara-

phin, 1971). Electron energy levels can be described by "band diagrams", where five lower bands originate from fully occupied d-levels, while one band near the Fermi level corresponds to the s-electron level. The shapes of the band diagrams and the Fermi surfaces of palladium, platinum, rhodium, iridium, and the hexagonally crystallizing ruthenium and osmium show increasing complexity in the direction of decreasing atomic number. Electrical Conductivity In a pure and perfect metal crystal, electrons would move freely through the lattice at absolute zero, resulting in zero electrical resistance. In real crystals collisions occur due to thermal vibrations of the lattice (phonons) - £ ph , impurity atoms or ions - Q{ , and crystal defects (dislocations, grain boundaries, etc.) - Qd, giving rise to increasing electrical resistance Q* = Qph + Qi + Qd

(9-1)

At low impurity concentrations, gph is independent of the number of defects, and gd is independent of the temperature (Matthiesen's rule). Considerable differences - of an order of magnitude of 6 between silver and palladium - exist for the resistivities of noble metals (Table 9-7). The low resistivity values of silver and gold (also copper) arise from their high charge

Figure 9-5. Fermi surfaces of (a) silver and (b) gold (Yonemitsu, 1991).

487

9.2 The Elements

Table 9-7. Electrical resistivities of NMs at 0°C in uX2 cm (Diehl, 1995). Noble method

Ru

Rh

Pd

Ag

Os

Ir

Pt

Au

Electrical resistivity Measurement error +

6.7

4.34 0.01

9.725 0.015

1.465 0.005

8.2 0.1

4.7 0.05

9.825 0.025

2.03 0.01

Table 9.8. Residual resistivity ratios (RRR) of NMs (Diehl, 1995). Noble method RRR a

Ru

Rh

Pd

Ag

o

IR

PT

Au

25 000

570

570

2100

400

85

5000

300

carrier density, given by their monovalent structures with one electron par atom in the outer s-band and the small amplitude of their lattice vibrations. The higher resistivity values of palladium and platinum are explained by their lower carrier density caused by the lacking on one d-electron and the increased scattering of electrons by phonons, caused by the density of d-states at the Fermi level. Their electronic structure is also responsible for the special T2 dependence of the resistivity at higher temperatures (Tanuma, 1991). Generally, Qph decreases with decreasing temperature, and finally becomes negligible at very low temperatures. The remaining amount of ge depends on the content of foreign atoms and lattice defects. The ratio of the resistivity measured at 273 K ( = £273) to £0 (the extrapolated resistivity for 0 K) is called the residual resistance ratio (RRR). It represents a measure of the purity of a metal. It is used to control the purity of platinum metal. For practical reasons, g4<2 is used, which can be measured at the temperature of liquid helium. Values measured for pure noble metals are listed in Table 9-8. In low-alloyed noble metals, the resistivity increases in proportion to the atomic concentration of the alloyed metal. With the addition of "B"-

metals (see Fig. 9-1) to silver and gold, the resistance of the alloy increases roughly with the square of the distance of the position of the B-metal from silver or gold, respectively in the periodic table (Table 99; Linde, 1932). The specific electrical resistance of solid solutions of noble metals and their temperature dependence are important parameters for selecting special alloy compositions for electrotechnical applications. The temperature coefficient of Table 9-9. Increase of atomic electrical resistivity AQ/C (in \xQ cm/at%) of silver and gold with alloying elements (Raub, 1940, 3). Alloying element

Al

Ag Au

1.87 1.9

Si

Alloying element

Ni

Cu

Zn

Ga

Ge

Ag Au

0.8

0.07 0.4

0.63 0.94

2.3 2.2

5 5.5

Alloying element

Pd

Ag

Cd

In

Sn

Ag Au

0.44 0.41

0.35

0.35 0.60

1.6 1.4

4.5 3.5

Alloying element

Pt

Au

Hg

Tl

Pb

Ag Au

1.5 1.0

0.36

2.3 0.8 0.41 14.4

4.6 3.9

488

9 Noble Metals and Their Alloys

resistance (TCR) aT is defined by aT=lx(Sr-QTl

(9-2)

Table 9-10 gives some selected values. Superconductivity has been found for some noble metals below the transition temperature Tc: ruthenium (Tc = 0.49 K), rhodium (Tc = 0.9 K), osmium (7; = 0.71 K), and iridium (Tc = 0.14 K). Higher transition temperatures are shown by some binary and ternary noble metal alloys (see Table 9-11; Khan, 1984; Raub, 1984; BenTable 9-10. Specific electrical resistance TCR (a r x 103) of NM solid solutions (Diehl, 1995). Solute content (at.c

Base / solute metal/ metal Ag/Au

025

axlO 3

Ag/Pd

025

axlO 3 Ag/Pt

025

80

60

40

20

8.3 0.93 11 0.58 33.0

11.0 0.83 22 0.40 60

11.0 0.84 41

8.1 1.1 34 0.75 35

9.8 0.88 28 0.28 21 3.8

17 30 0.61 0.45 44 37 0.26 0.82 37 57 1.8 -<0.1

46

axlO 3

Au/Pd

025

axlO 3

Au/Pt

025

axlO 3 Rh/Ni

025

axlO3

26 1.2 34 0.8 50 3

Table 9-11. Superconducting noble metal alloys (Benner etal., 1991,1). Alloy Nb 3 Au Nb 3 Pt ZrRuP Tb 3 Os 4 Ge 13 ZRuAs YOs4Gel0 ScIr4Si10 ErRhB 4

Transition temperature (K) 10.9 10.6 12.3 14.1 11.9 8.6 8.4 8.5

ner et al. 1991,1). In Pt-Mo alloys, a maximum 7^ of 7.5 K occurs at a composition of 30 at.% Pt, with a partially ordered superlattice phase of the hexagonal DO 1 9 type (Flukiger et al., 1972); (see also Sect. 9.3.2.3). Superconductivity occurs preferably in materials with strong electron/lattice interactions, which normally correspond to high electrical resistivities (Bardeen et al., 1957). The superconductivity is not sensitive to nonmagnetic impurities, but is strongly influenced by ferromagnetic impurities. These lower the transition temperature by interaction with the electron spin, and destroy the superconductivity forming electron pairs. Among the ternary noble metal alloys, there is a special group, called magnetic superconductors, which exhibit coexistence of a superconducting phase and magnetic order, which can be seen in the temperature dependence of the electrical resistence and the magnetic susceptibility. Below a critical temperature Tcl, superconductivity and diamagnetism coexist over a particular temperature range. On further lowering of the temperature to below a second critical temperature Tc2, the electrical resistance rises again and the magnetic moments become ferromagnetically ordered. One example of a compound that exhibits such behavior is ErRh 4 B 4 (Fertig et al. 1977; Maple etal., 1980). Other combinations show a relation between antiferromagnetism and superconductivity. Thin sandwich films of gold/ chromium/gold with a chromium layer thickness of 1.5 nm are superconducting. While the antiferromagnetic chromium has in bulk form a b.c.c. crystal structure, it changes in the thin film form to an f.c.c. crystal structure. This transformation is supposed to be responsible for the superconductivity of these sandwich films, where

9.2 The Elements

the thin chromium layers form the superconducting phase and the gold layers merely act as nonsuperconducting, mechanical supports to hold the chromium layers between them (Brodsky et al. 1982). Further superconducting noble metal alloys have been found in the alloy systems Y-Pd-B-C and Y-Pd-B with Tc value of ^23 K (Cava etal, 1994). In the system SrRuO, an antiferromagnetic phase with a perovskite-type structure becomes superconducting at T c «1 K (Maeno et al., 1994). While in the compound LaBa 2 Cu 3 O 7 , the superconductivity is ascribed to the layer structures of CuO; the superconductivity of SrRuO is believed to be due to the Ru-O planes and may originate from a different mechanism. All PGMs catalyze the formation of YBA 2 Cu 3 O 7 by lowering the reaction temperature and accelerating the reaction by a factor of six. The catalyzed reactions are sometimes combined with the formation of semiconducting noble metal compounds (Shul'ga et al., 1993). As well as in the solid, metallic state, electrical conductivity is also found in a special group of complex chemical PGM compounds. Platinum and its allied 8group metals form organometallic complex ions with ligands coordinated in tetrahedral symmetry around the platinum atoms. They can be crystallized out of solutions with charge-compensating ions, forming linear-shaped chain molecules. These crystals represent bundles of chain molecules. In the presence of halogen ions, which act as strong electron acceptors, these chains are electrically conducting along the chain axes and are quasi onedimensional metallic conductors. At temperatures above 250 °C, the conductivity of a metallic character, at lower temperatures it decreases with decreasing temperature with a gradual transition to an insulating state ("Peierls transition"; Peierls, 1955).

489

Two chain types with different bond structures are distinguished: 1. Chains formed of Me-Me along the axes (Me = metal): This group is formed with platinum, rhodium, and iridium. A representative type of the numerous compounds is a coppery red compound of the composition K 2 [Pt(CN)4] • Br0>3 • x H 2 O ("KCP" and "KCP family"; Minot and Perlstein, 1971). The distance between the platinum ions in the chain is 0.288 nm, which is close to the atomic distance of 0.278 nm in platinum metal with its closepacked f.c.c. structure. In the absence of bromine ions, the compound is an insulator. Electrical conductivity along the Me-Me bonds becomes possible by the transfer of electrons from the metal to the bromine, which acts as strong electron acceptor (Kroogmann and Hausen, 1968). 2. Chains consisting of alternating Me and halogen bonds: These halogen-bridged, mixed-valence chain crystals are formed by platinum and palladium. A representative type of the numerous compounds is "Wolffram's red salt" [Pt IV (C 2 H 5 NH 2 ) 4 ] [Pt n (C 2 H 5 NH 2 ) 4 Cl 2 ]Cl 4 • 4H 2 O (WRS). In this compound type, electron transfer between the metal ions becomes possible via p-orbitalis of the halogen ions. In compounds with chlorine ions, the interatomic distances between the halogen and the metal ions are smaller for the metal ions of higher valency and larger for metal ions of lower valency. By changing the halogen from the "light" chlorine ion to halogens of higher atomic number, the electron cloud becomes larger with increasing atomic number and electron transfer between the metal ions becomes easier. With increasing atomic number, the position of the halogen ions also comes closer to the midpoint between the metal ions. With iodine ions the structure approaches an idealized, one-

490

9 Noble Metals and Their Alloys

dimensional metallic conductor, with a valency structure which corresponds to an average valency for all the metal ions of 3 + instead of the alternating distribution of 2 + and 4 + found with the lighter halogen bridge ions (Tanino, 1982; Kobayashi, 1983). Crystal Structures The crystal structure of the transition metal change with increasing atomic numbers from the more "open" body-centered cubic structures to the close-packed structures of highest symmetry and highest coordination numbers. Ruthenium and osmium have hexagonal close-packed (h.c.p.) crystal structures, and rhodium, palladium, silver, iridium, platinum, and gold have face-centered cubic (f.c.c.) crystal structures. There are octahedral and tetrahedral holes in the f.c.c and the h.c.p. structures which can be occupied by interstitial atoms having maxium atomic radii of 0.4 r or 0.225 r, respectively (r = atomic radius of the noble metals, if these are considered as close-packed spheres in these lattices). Densities All the noble metals have densities above 10g/cm 3 , the heavy noble metals above 19 g/cm3 (Table 9-6). The density of the transition metals increases due to the completion of their 3d-, 4d-, and 5d-electron levels with a constant outer s2-level, with increasing nuclear charge. The highest density values are reached in the 2nd long period at ruthenium and in the 3rd long period at iridium. Further completion of the lower d-shells results in a slight and then stronger decrease in the density, accompanied by increasing atomic diameters. The much higher densities of the noble metals of the 3rd long period compared to those of their analogs of the

2nd long period are caused by the full, occupied 4d-levels, which are completed in the rare-earth group (lanthanides) and cause a marked increase of the elements' densities with only a small increase in the atomic diameters. Mechanical Properties Basic mechanical property values are given in Table 9-6. Noble metals exhibit a wide range of ductility. Silver, gold, palladium and platinum have low strength and can easily be cold worked. The work-hardening is not very marked. Gold can be cold deformed to more than 99% of its crosssection without intermediate annealing. High-purity gold wires can be cold drawn down to diameters of 10 jim. Very thin, beaten gold leaf is hammered to as thin as 0.1 |im; the minimum thickness of mechanically deformed silver foils is « 0.2 jim. The recrystallization temperatures of 0.28 Tm (Tm = melting temperature in kelvin) for silver and gold, ^0.31 Tm for platinum and ^0.41 Tm for palladium are comparatively low; they decrease with increasing degree of cold deformation (Jonsson, 1995). Rhodium and iridium have much higher strengths and have only limited workability at room temperature. They can be easily deformed at higher temperatures (rhodium >200°C, iridium >600°C), accompanied by strong work hardening. The recrystallization temperatures are, with ^0.43 Tm for rhodium and ^0.55 Tm for iridium, comparably high. Ruthenium and osmium exhibit high strength. They can only be deformed at high temperatures, and only to a limited degree. Work hardening for these metals is stronger than for rhodium and iridium. Ruthenium, which has been deoxidized by melting with rare earths, can be hot rolled. A more sophisticated method of deforming

491

9.2 The Elements

this metal is the composite approach by sintering ruthenium powder in the presence of a liquid phase that solidifies to a ductile matrix which surrounds the particles. Palladium-gold alloys are suitable for this process, resulting in duplex alloys of reasonable ductility (Savage and Tracey, 1971). Ruthenium, with ^0.57 Tm, has the highest recrystallization temperature. The decrease in the recrystallization temperature with increasing nuclear charge of the noble metals in both periods runs parallel with decreasing melting points, decreasing hardness, and decreasing brittleness in the same direction, indicating that the interatomic bond strength becomes weaker on completion of the lower d-shells above the d8 -configuration. The thermal expansion and the ductility increase in the same direction. Gold has the highest ductility. Figure 9-6 shows the increase of hardness caused by cold deformation. The rate of work hardening increases with increasing modulus of elasticity. There are two separate curves for the cubic and hexagonal structures (Fig. 9-7) (Darling, 1973, 1). The deformation behavior can be related to the ratio of the bulk modulus to the rigidity modulus (Table 9-12). Lower values correspond to the harder and more brittle metals (Pugh, 1954). The mechanical properties of the elements depend on the temperature (Table 9-13). The elastic modulus of the single crystals of the f.c.c. noble metals shows considerably different values in the [100], [110], and [111] directions, causing anisotropic elastic properties of the single crystals (Table 9-14), and in polycrystals with textures. Thermoelectric Properties The thermoelectric properties of noble metals are of great practical importance.

'•5 600 r

Re

Ir

0 —

10 20 30 40 50 Percentage reduction in thickness

Figure 9-6. Work hardening of the PGMs (Darling, 1973, 1).

1400

olr

;200

! 100

g ~Z

V

A u

5000

10000

15000

20000 25000

Modulus of rigidity (kg/mm2)

(V

Figure 9-7. Work hardening of cubic and hexagonal metals (Darling, 1973, 1).

Thermocouples made from NM alloys are used for high temperatures measurements because of their refractoriness and oxidation resistance. The thermoelectric voltage is defined as the voltage E A p l a t i n u m in mV, which a wire made from material A develops in conjunction with a wire made from physically pure platinum at a temperature difference between the two measurement points. The thermoelectric voltage is

492

9 Noble Metals and Their Alloys

Table 9-12. Mechanical properties of PGMs (Darling, 1973, 7). Metal

Os Ir Re Ru W Rh Pd Pt Au

Crystal structure

Young's modulus (GPa)

Modulus of rigidity, G (GPa)

Bulk modulus, K (GPa)

Poisson ratio

K/G

h.c.p. f.c.c. h.c.p. h.c.p. b.c.c. f.c.c. f.c.c. f.c.c. f.c.c.

560 538 472 430 396 386.4 128.3 174 80.2

220 214 180 172 151.4 153 46.1 62.2 28.2

380 378 340 292 318.6 280.1 190.9 280.9 174.6

0.25 0.26 0.26 0.25 0.29 0.26 0.39 0.39 0.42

1.73 1.76 1.89 1.71 2.11 1.83 4.13 4.52 6.18

Table 9-13. Temperature dependence of the modulus of elasticity (Jonsson, 1995). Metal

7TQ

E (GPa)

Rh

20 900

386 291

Pd

20 800

124 91

Ag

20 900

82 40

Ir

20 1000

538 434

Pt

20 1000

173 115

Table 9-14. is-moduli in different crystal directions (Jonsson, 1995). Direction

E (GPa)

Ag

[100] [110] [111]

44 82 115

Au

[100] [110] [111]

42 81 114

Pd

[100] [110] [111]

65 129 187

[110] [111]

480 660

Noble metal

Ir

positive if on the hot junction the current flows from the platinum to material A. It is a material characteristic and defines the basic thermoelectric behavior of different materials. Table 9-15 gives values of the thermoelectric voltage of pure noble metals against platinum at different temperatures. The thermoelectric voltage is changed by alloying elements (Table 9-16). Compositions and thermoelectric voltages of noble metal containing thermocouples are given in Sect. 9.5.2.4. Minute quantities of magnetic impurities, particularly iron, down to the parts per billion level, influence the low temperature electrical transport properties of gold ("Kondo effect"). They cause a minimum in the electrical resistivity, accompanied by anomalies in the thermoelectric power, magnetostriction, and magnetic susceptibility (Kondo, 1969). These changes can be used to determine the impurity concentrations below lOppm. Measurements of thermoelectric power can therefore be used to determine the content of iron in gold down to ^0.01 ppm (Kopp, 1976). Magnetic Properties The magnetic properties of the elements are described and measured by the mass

493

9.2 The Elements

Table 9-15. Thermoelectric voltages of noble metals3 (Diehl, 1995).

7TC)

Ru

Rh

Pd

Ag

- 50 + 100 + 200 + 800

0.684 1.600 9.519

-0.2 0.70 1.61 10.16

+ 0.2 -0.570 -1.23 -7.98

-0.1 0.74 1.77 13.36

a

Ir -0.2 0.66 1.525 9.246

Au -0.1 0.77 1.834 12.288

E,. Pt in mV.

Table 9-16. Thermoelectric voltages of NM alloys3 (Diehl, 1995). Alloy Base

Addition

Rh Pd

Ir Ag

Pd

Au

Pd

Pt

Pd

Ni

Ag

Au

3

Composition (added alloy component wt.%)

T(°Q

1000 100 1000 100 1000 1300 100 1000 1300 100 1000 100 700

10

30

60

80

15.15 -1.1 -23.5 -1.0 -14.5 -22.0 0.32 -2.1 -5.2 -0.8 -9.4 0.56 8.6

17.40 -2.4 -44.4 -1.7 -24.1 -34.0 0.83 7.8 8.0 -1.47 -9.4 0.44 7.0

18.25 -1.2 -26.2 -3.5 -42.7 -54.0 0.63 8.9 11.8 -1.75 -10.9 0.42 6.8

17.05 -0.1 -1.1 -0.5 -10.5 -14.0 0.37 6.5 8.6 -1.73 -11.9 0.49 7.3

Ein mV.

susceptibility #m, which is composed of the diamagnetic portion #Dia and the paramagnetic portion #Para according to Zm ~~ ZDia ' XPara

(9-3)

with

_c APara

rp

(9-4)

Curie's law and (9-5) ~ 3R where M is the magnetic moment per mole and R the gas constant. The magnetism of the noble metals is very weak. In the metallic state, silver and L

gold (also copper) are diamagnetic, while the platinum group metals are paramagnetic. Table 9-17 gives the susceptibilities at 0°C and 20 °C. The diamagnetic susceptibilities of silver and gold are independent of the temperature between zero and the melting point. The paramagnetic susceptibilities of palladium and platinum decrease with increasing temperature, while in contrast the susceptibilities of rhodium, iridium, ruthenium and osmium increase with increasing temperature. Measuring the magnetic moments is used to determine the oxidation state of elements in alloys and compounds. Because of the very narrow energy levels between the outer s-shells and the lower d-electron levels of the last mem-

494

9 Noble Metals and Their Alloys

Table 9-17. Susceptibilities of NMs (Diehl, 1995). X(20°C) Z(0K) ( 1 0 - 9 m 3 k g - 1 ) (10" 9 m 3 kg- 1 )

Metal

4.85 11.56 82.68 0.60 1.29 13.32 -2.5 -1.7

Ru Rh Pd Os Ir Pt Ag Au

5.15 12.44 63.89 0.69 1.51 12.19 -2.5 -1.76

where / is the sample length and st the specific magnetostriction (Table 9-18). Based on this and in combination with ordering processes, alloys in the systems PtFe and Pd-Fe show around the Fe 3 Pt and Fe 3 Pd stoichiometry (like Fe-Ni alloys, e.g., Ni-Fe36, "Invar") with zero or negative coefficients of thermal expansion for similar magnetic properties (Kussman and Von Rittberg, 1950; Kussman and Jesson, 1962). Optical Properties

bers of the transition metals with nearly fully occupied d-levels, interactions occur by completion of the lOd-shell. Paramagnetic palladium dissolves in the diamagnetic copper, forming a diamagnetic alloy. Palladium does not contribute conducting electrons in these alloys. Obviously, the partial dissociation of metallic palladium according to Pd r

:Pdp + a r a

(9-6)

which is responsible for its paramagnetism, is suppressed by the high electron-gas concentration of the metallic copper. The ordered phases in this composition range (see Sect. 9.3.2.3) have higher diamagnetic susceptibilities than the disordered solid-solution phase. Only disordered Cu-Pd alloys with palladium contents of above 50 at.% are paramagnetic. Similar behavior is found when hydrogen is dissolved in palladium, where the paramagnetism of palladium decreases down to zero at the composition Pd-H 0 66 leaving a pure, diamagnetic Pd-H alloy (Lewis, 1967). In these "alloys", the hydrogen occupies octahedral interstitial sites in the palladium lattice. PGMs show magnetostriction under the action of a magnetic field H Al

T

(9-7)

The optical properties are closely related to the interactions of photons with s- and d-electrons. Optical constants give information obout electronic band structures, the participation of free and bonded electrons, about phase formation, ordering processes, and lattice defects. Colors are due to selective light absorption by interband transitions of electrons caused by the absorption of photons. Photons can be absorbed by exciting an electron of a filled band just below the Fermi surface. The energy difference between this band and the Fermi surface determines the minimum energy of the photon which can be absorbed by the interband transition.

Table 9-18. Magnetostriction of NMs (Diehl, 1995). Metal (10Ru Rh Pd Ir Pt Rh,O.5O^rO.5O Rh•0.50*^*0.50 Pdo.67Pto.33 Pdo.33Pto.67

••A-)

-1.4 11 -39.4 3.8 -32 9.5 27 13.4 -17.4 -79

9.2 The Elements

495

1.0 r 0.8 >> " | 0.6 \ 0.4

Figure 9-8. Reflectivity of noble metals in the wavelength range 0.2-5 |nni for normally incident light (Volcker, 1995).

0.2

0 0

0.5

1.0

1.5

2.0

2.5

3.0

3.5

4.0

4.5

5.0

Wavelength (pm)

Figure 9-9. Reflectivity of noble metals in visible light (Volcker, 1995).

0.4 520

580

640

Wavelength (nm)

The reflectivity and color of the noble metal are of special importance for jewelry and dentistry alloys, and for optical and heat-reflecting mirrors. Emissivity values are needed for the calibration to temperature measurements by pyrometry. The reflectivity is defined by the ratio of the intensity of the reflected to the incident light (Hummel 1971, 1) r

(9-8)

opt

For vertically incident light, the reflectivity is connected to the refraction index n and the extinction coefficient x by the relation opt

~

(

'

Figures 9-8 and 9-9 show the reflectivity of bulk noble metals in the wavelength

range from 0.2 to 5 ^m, and for visible light. The high reflectivities of the noble metals are in accordance with the HagenRubens relation (Hagen and Rubens, 1903) (9-10) (where v is the optical frequency of light and o-e the electrical conductivity), which states that at low frequencies metals with high electrical conductivities are also good optical reflectors. According to the Hagen-Rubens relation, silver has the highest reflectivity in the visible light range and in the infrared. Silver coatings are therefore applied for optical mirrors, reflectors, and for transparent, heat-reflecting architectural glasses for windows, doors, etc. For optical mirrors and reflectors, the surfaces are protected against corrosion by

496

9 Noble Metals and Their Alloys

coating with thin layers of rhodium; architectural glasses are protected by optically transparent oxide coatings, which act at the same times as antireflection layers. The reflectivity decreases very slowly with decreasing wavelength in the infrared and in the visible range, but shows a steep drop to a low minimum of only a few percent in the ultraviolet branch at ^314 nm. This is ascribed to an interband transition by the obsorption of photons with a critical threshold energy of ^4eV, which brings one d-electron up to the Fermi surface (Hummel, 1971, 2). The reflectivity of gold shows a marked fall at ^550 nm in the visible range, with a minimum of r opt «0.25 in the near ultraviolet. Because d-bands are positioned at higher energy levels, interband transitions occur at smaller energies («2.5 eV) in the visible range. The reflected light contains all wavelengths above 560 nm, which gives the typical gold-yellow color. Because of gold's high reflectivity in the infrared range, gold films are used in radiant heating and drying devices, for thermal-barrier windows of large buildings, and as reflective coatings to protect space vehicles and space suits against excessive solar radiation. Ultrafine particles of gold ("goldblack") show selective light absorption. Reflectivity is less than 0.01 in the visible region and less than 0.1 in the infrared. Gold-black therefore forms an excellent light-receptive surface coating for radiation detectors, and is also suitable as a selective absorption film for the capture of solar energy (McKenzie, 1978). The reflectivities of the platinum group metals are lower. The characteristic fall with decreasing wavelength in the visible range is smallest for rhodium which has a neutral color but a reflectivity ^ 2 0 % lower than that of silver. Because of its high corrosion resistance, thin films of rhodium

are used to protect silver surfaces against corrosion. Osmium and iridium are efficient reflectors in the wavelength region below 100 nm (Hass and Hunter, 1974) and are therefore used for mirror coatings for vacuum UV (ultraviolet) spectroscopic to study celestial phenomena which cannot be found with aluminum-based mirrors. Iridium-coated mirrors have been used in large aperture telescopes of space rockets for exploration tasks in enhanced space astronomy (Herzig and Spencer, 1982; Herzig, 1983). The optical transmission of thin films depends on the wave-length and the film thickness. Thin gold films appear green in transmission. The penetration of photons in the metal surfaces is in the range of 10~ 6 cm. Above this thickness, transmission decreases very rapidly. Colored noble metal alloys are found with different structures and compositions. Table 9-19 gives some examples. Alloys of Table 9-19. Colored noble metal alloys (Hummel, 1971, 4). Alloy

Color

Remarks

Ag-Zn (/?-phase) Ag-Au(70) Al2Au KAu 2 Au-Zn-Cu-Ag Auln 2

rose green-yellow violet violet green blue

Zintl phases Li2AgAl Li2AgGa Li2AgIn Li2AgTl Li2AuTl

yellow-rose yellowish gold-yellow violet-rose green-yellow

VEC VEC VEC VEC VEC

1.5 1.5 1.5 1.5 1.5

rose-violet rose-violet violet blue-violet violet

VEC VEC VEC VEC VEC

1.75 1.75 1.75 1.75 1.75

Li2AgSi Li2AgGe Li2AgSn Li2AgPb Li2AuPb

VEC = Valence electron concentration.

497

9.3 Alloys Table 9-20. The emissivity of noble metals at 1000 °Ca (Smithells, 1977). Metal Emissivity

Ag

Au

Rh

Pt

Ir

Pd

Ru

Os

0.0551

0.16

0.22

0.29

0.36

0.37

0.42

0.52

at wavelength 0.65 jxm,

b

at 800 °C.

the composition Li2NMX (where X = metal of group 3 or 4) have the structure of Zintl phases if they are colored (Hummel, 1971,3). The emissivity is defined as e= 1

' opt

(9-11)

and depends on the material, the wavelength of the radiation, the temperature, and the surface conditions. It is important to calibrate deviations against the emission valus of a black body in pyrometrical temperature measurements. Table 9-20 gives some values (Smithells and Brandes 1977).

9.3 Alloys 9.3.1 General Remarks Noble metals form numerous alloys and intermetallic compounds with each other and many B- and A-group metals. These alloys have in most cases f.c.c, h.c.p., or b.c.c. crystal structures, but there are also a great variety of different structure types for the intermetallic compounds. Many of these are stable over an extended composition range and also represent metallic phases rather than intermetallic compounds of definite stoichiometric composition. 9.3.2 Phase Formation in Noble Metal Alloy Systems 9.3.2.1 Primary Solid Solutions Solid solutions in noble metal alloy systems with metals are generally of the sub-

stitutional type, where the solute atoms replace some of the solvent atoms on their lattice positions. Palladium and platinum form interstitial solid solutions with small atoms like hydrogen and carbon. In the solid-solution range of some systems, ordering reactions occur with the formation of superlattices of the same or similar lattice types. These transformations are accompanied by changes of the mechanical and physical properties. They find practical applications in the hardening of noble metal alloys and in the production of magnetic alloys. Generally, substitutional solid solutions are only formed if the components have the same or similar crystal structures, and if the atomic sizes of the solvent and the solute differ by not more than « 1 5 % ("size factor"; Hume-Rothery etal, 1934). Continuous solid solutions are formed between most noble metals, and between gold, palladium, and platinum, and the group homologs of the first long period: iron, cobalt, and nickel. Figure 9-10 shows these binary alloy systems and Table 9-21 gives the corresponding size factors in terms of the percentage difference in the atomic diameters. Some binary systems with continuous solid-solution ranges immediately beneath the solidus line (Au-Ni and Au-Pt) form miscibility gaps with temperature-dependent phase boundaries at lower temperatures. The tendency for homogeneous alloys to split up into two mutually saturated solid solutions has been thought to correlate with the difference of the melting temperatures of the alloy con-

498

9 Noble Metals and Their Alloys

Figure 9-10. Binary systems forming continuous solid solutions, indicated by connecting lines.

causes a strong increase in the miscibility gap, leading to complete phase separation. Systematic relations exist for the alloying behavior of silver and gold with Bgroup metals. With constituents of favorable size factors, solubility limitations are give by the ratio of the number of valency electrons per atom, the "electron concentration" (e/a), calculated by (Hume-Rothery, 1956, 1) e

Table 9-21. Size factors for systems with continuous solid solutions. Alloy system Co-Rh Co-Ir Co-Pd Co-Pt Ni-Pd Ni-Pt Cu-Pd Cu-Au Rh-Ir Rh-Pd Ir-Pt Pd-Pt Pd-Au Pd-Ag Au-Ag

Atomic diameter difference 7.3 8.3 9.8 9.3 10.4 10.2 7.2 11.4 0.9 2.3 1.2 0.2 4.8 5.0 0.2

stituents (Raub, 1959). The critical temperature of the miscibility gap of the alloys Pt-Ir, Pt-Rh, Pt-Au, and Au-Ni increases with increasing difference of the melting temperatures of their components. In the system Pt-Au, the miscibility gap is already near to the solidus. On the basis of the above consideration, a qualitative interpretation may be given for the formation of the peritectic phase diagram of Pt-Ag, where the somewhat greater difference in the components' melting temperatures

(9-12)

100

a

where V = valency of the solvent, and v and x are, respectively, the valence and the atomic percentage of the solute. This amounts for alloys of silver and gold with B-group elements to ^1.4 for silver and ^1.3 for gold. These values decrease with increasing valency of the solute (Table 9-22). Increasing valency of the solute results in a more restricted solid solution. Alloys of identical, equivalent composition (i.e., atomic percentage of the solute multiplied by its valency), have identical liquidus points within the limits of accuracy (Hume-Rothery and Reynolds, 1937). In accordance with the general principle of mutual solubilities of components of different valency in binary alloys (HumeRothery, 1956, 2), the solubilities for solutes with a valency higher than one are always higher on the side of silver and gold (also copper). Table 9-23 gives some values for the maximum solid solubilities of B-

Table9-22. Electron-concentration values for gold and silver alloys (Hume-Rothery, 1956, 6). Silver alloys, e/a Ag-Cd Ag-In Ag-Sn Ag-Sb

1.428 1.404 1.335 1.29

Gold alloys, e/a Au-Cd Au-In Au-Sn Au-Sb

1.33 1.255 1.205 1.045

Ae/a 0.098 0.149 0.130 0.245

499

9.3 Alloys

Table 9-23, Solid solubilities of B-elements in silver and gold, and of silver and gold in these base metals (Hume-Rothery, 1956, 7). Ag

Au

Maximum ssolubility

Maximum :solubility

Solute metal

Mg Zn Cd Hg Al Ga In Tl Si Ge Sn Pb As Sb Bi

Solubility (at.%)

Temperature (°C)

Ag (at.%)

Temperature (°C)

Solubility (at.%)

Temperature (°C)

Au (at.%)

29.3 40.2 42.2 37.3 19.9 18.7 20 13.2 _ 9.6 12.5 2.8 8.8 7.27 0.8

759 258 300 276 448 611 693 550

4 5 7 24 <3 _ _ -

471 431 343

32.5 31 32.5 2.6 16 12 12.6 <1 _ 3.2 6.8 1.1 -

827 642 625 420 545 455 647 800

_ 5 3.5 _ _ 0.2 — -

651 724 600 595 702 259

566 25

metals in silver and gold, and the maximum solid solubilities of silver and gold in these B-metals. In each case, a favorable size factor is the necessary precondition for any marked solid solubility. The actual size factor may be further influenced by the degree of ioniziation of the solute atoms. The solid solubilities of A-group elements in silver and gold cannot be coordinated to definite e/a values. The solubilities of the transition elements of the first long period are remarkably higher than those of the second and third (Table 9-24). This is ascribed to the progressively increasing participation of d-electrons of the metals with a b.c.c. structure, forming directional bonds with increasing atomic number (Raynor, 1976, 1). The solid solubility in ternary systems is also determined by the tendency to maintain a constant, limited electron :atom ratio. Differences in valency and atomic sizes causing lattice distortion

356 498

550

Temperature (°C) 423 309

217

Table 9-24. Solid solubilities of A-elements in gold (Raynor, 1976, 4). Solute

Maximum solubility (at.% solute)

Temperature (°Q

4A

Ti Zr

13 7.25

1123 1065

5A

V Ta

59 11.3

1400 1000

6A

Cr Mo

49.7 1.25

1160 1054

7A

Mn Re

30.8 0.1

960 1000

8a

Fe Ru

75 0.5

1168 950

8b

Co Rh

23.5 0.56

996 900

8c

Ni Pd

Group of the periodic table

100 100

500

9 Noble Metals and Their Alloys

of the solvent effect changes in the solid solubility in the solvent-rich corner. The solid solubility is lowered in the system Ag-Mg-Sn, while increased solid solubility is found in the systems Ag-Zn-Sb and AgCd-Cu. 9.3.2.2 Intermetallic Compounds

Table 9-25. NM 3/2-compounds (Raynor, 1976, 2). Phase range

Crystal structure type

Structure

Alloy system

P

W

A2 cubic OjJ-lm3m

Au-Al(HT) a Ag-Al

F

CsCl

B2 cubic Oi-Pm3m B19 orthorhombic Dfh-Pmcm

Au-Mg Au-Zn Au-Cd

n

|5-Mn

A13 cubic O6-P423 O7-P413

Au-Al(LT) b Au-Al

Q

Mg

A3 hexagonal D6h-P63mmc

Au-Al Au-In Au-Sn Ag-Al Ag-Sn

el a Compounds (Hume-Rothery Phases) Copper, silver, and gold form intermetallic compounds with B-metals with identical or similar crystal structures at compositions corresponding to e\a values of 3/2, 21/13, and 7/4 ("Hume-Rothery phases"). Most of these phases extend over a certain homogeneity range. The relation between these el a values and certain crystal structures can be explained basically by the band theory of free electrons. The observed values of el a are near the electron concentration at which the Fermi sphere meets the boundary of the first Brillouin zone of the corresponding crystal structures: 1.36 for the f.c.c. (a), 1.48 for the b.c.c. (p, P'), 1.54 for the complex cubic (y), and 1.69 for the h.c.p. (e) phase (Kittel, 1986). The 3/2 electron compounds form different structures, which are determined by the influence of solute valency, temperature, size factor, and differences in the electronegativities (Table 9-25). The formation of ordered p'-phases is favored by an increased difference in the electronegativities of the components, leading to the occupation of neighboring lattice sites by unlike atoms. P'-phases are found in the binary gold alloy systems, Au-Zn, Au-Cd, and Au-Mg, whereas in binary silver systems the P'-phase is only present in Ag-Mg (Raynor, 1976, 2). The y-brass structure is formed in binary noble metal alloys at compositions which approximate to an el a value of 21/

a

HT = High-temperature modification. b LT = Lowtemperature modification.

13 (Table 9-26). The complex cubic structure contains 52 atoms per unit cell, which could theoretically, if all electron levels are occupied, take 90 electrons in the electronic energy band. Its homogeneity range reaches from e/a = 1.54-1.7, corresponding to a maximum occupation number of 88 electrons per unit cell. The conditions for phase stability depend on the formation of lattice distortions and electrochemical interactions between the components. Both can give rise to displacement of the homogeneity range to higher B-contents. This is found in the Au-Zn system, where the y-phase extends beyond the e/a limit of 1.7. In this case, the y-phase is stabilized by the formation of a defect structure by omitting the corresponding number of atoms to maintain the maximum number of 88 electrons per unit cell. No y-phases are formed in the systems silver-aluminum, silver-magnesium, and gold-magnesium. The very small lattice distortion in the sil-

501

9.3 Alloys

Table 9-26. e/a compounds (Hume-Rothery, 1956, 8). Electron: atom ratio = 3 :2

Electron:

Electron:

Bodv-centered cubic structure

Complex cubic ("/^-manganese") structure

Close-packed hexagonal structure

atom ratio — Zi .13

atom ratio — 7:4

"y-brass" structure

Close-packed hexagonal structure

CuBe CuZn Cu3Al Cu3Ga Cu3ln Cu5Si Cu5Sn AgMg AgZn AgCd Ag3Al Ag3ln AuMg AuZn AuCd FeAl CoAl NiAl Niln Pdln

Cu5Si AgHg Ag3Al Au3Al CoZn3

Cu 3 Ga Cu5Ge AgZn AgCd Ag3Al Ag3Ga Ag3ln Ag5Sn Ag7Sb Au 3 ln Au5Sn

Cu 5 Zn 8 Cu 5 Cd 8 Cu 5 Hg 8 Cu 9 Al 4 Cu 9 Ga 4 Cu 9 ln 4 Cu 31 Si 8 Cu 31 Sn 8 Ag 5 Zn 8 Ag5Cd8 Ag 5 Hg 8 Ag 9 ln 4 Au 5 Zn 8 Au 5 Cd 8 Au 9 ln 4 Mn 5 Zn 21 Fe 5 Zn 21 Co 5 Zn 21 Ni 5 Be 21 Ni 5 Zn 21 Ni 5 Cd 21 Rh 5 Zn 21 Pd 5 Zn 21 Pt 5 Be 21 Pt 5 Zn 21 Na 31 Pb 8

CuZn3 CuCd3 Cu3Sn Cu3Ge Cu3Si AgZn3 AgCd3 Ag3Sn Ag5Al3 AuZn3 AuCd3 Au3Sn Au 5 Al 3

1 — ^ V V /»

V

V V l l t V l

VVf

ver-aluminum system stabilizes the closepacked hexagonal phase which forms a continuous field of solid solutions with the 7/4-electron compound. In the silver-magnesium and gold-magnesium systems, formation of the y-phases is suppressed by strong electrochemical interactions, causing displacement of the corresponding phase boundaries outside the valid e/arange (Hume-Rothery, 1956, 3; Raynor 1976, 3). The 7/4-type electron compounds of copper, silver, and gold alloys with B-subgroup metals have close-packed hexagonal structures.

Other Intermetallic Phases Besides these e/a determined structures, a huge number of other intermetallic phases exist in noble metal alloys, most of them are also of metallic character with predominantly cubic, hexagonal, or tetrahedral symmetry. Examples of typical structures are listed in Table 9-27. The CsCl structure is mainly present in compounds of gold and silver with rareearth metals and platinum group metals with 3A-transition metals. In the MoSi2-type structure (Cllb structure), formally regarded here as an

502

9 Noble Metals and Their Alloys

Table 9-27. Structure types of intermetallic phases. Type

Structure

NM phases

CsCl

B2 cubic

PdTi, PtLi, TiPt(H), PtZr, PtHf, PtMn, Au-RE a

MoSi2

C l l b tetragonal D^-i/mmm

AuZr2, Au2Zr, AuHf2, Au2Hf AuMn 2 , Au2Mn, Pt3Hf2, Pd2Hf, Au2Al, Au2Ti, Au-RE, Ag-RE, PdZr2, Pd2Zr, Pt2Si

C15 cubic

NaAu 2 , AuBe2,PbAu2, BiAu2, Pt2Li, Pt2Ca, Pt2Ba, Pt 2 -RE CaAg2, Zrlr 2 , ZrRu 2 , ZrOs2 MgZnAg, MgAuAg, MgAgAl

Laves MgCu2 MgZn2 NiAs

C14 hexagonal D*h-P63/mmc B8 hexagonal

PtB, Pt 3 ln 2 , PtSb PtFe, Pd 3 Pb 2 , Pd 5 Sb 3 , PdSb

C6 hexagonal MnP

B31 orthorhombic D^-Pcmm

CaF 2

Cl cubic Oh-Fm3m "isomorph" "anti-isomorph" (MgSn2-type) (3-cubic T^-F43m

1

AuGa, PdSi, PtSi, PdGe, IrGe, PtGe PdSn, RhSb, PdSb

PtSn2, Ptln 2 , AuAl2, AuGa 2 , Auln 2 , PtAl2 Ir 2 P, AgMgAs

RE = rare earth element.

AB2-type composition, noble metals can take either the A- or the B-position. The phases have different axial ratios when they are alloyed with 4A-group elements (zirconium, hafnium) or manganese. The metal in the predominating B-position determines the character of the bond. Low axial ratios, corresponding to a coordination number of 10, are formed with noble metals or B-subgroup elements in the Bposition. If the 4A-elements are placed in the B-position, they effect a stronger participation of directed, hybridized bonds. This results in a higher axial ratio with the coordination number 8 and stabilizes the b.c.c. structure. The Laves phases (C15 structure, Hume-Rothery, 1956, 4; Haasen, 1984, 1)

are formed by AB 2 compounds with an atomic diameter ratio of «1.2. The densely packed structures are predominantly determined by geometrical conditions. Depending on the ratio of their atomic diameters, A- and B-atoms are interchangeable. Noble metals are found in the cubic MgCu2 and the hexagonal MgZn 2 structure type in binary and, by replacing one B-component, also in ternary alloys (Massalski, 1983). Both structures are of high symmetry. A- and B-atoms are placed on two interpenetrating lattices with the larger atoms accommodated in holes in a skeleton of B-atoms in tetrahedral positions (Girgis, 1983). The p-tungsten phase (A 15 structure) occurs in stoichiometric phases with an

503

9.3 Alloys

AB 3 composition. The B-atoms occupy a b.c.c. lattice and have 12 A-neighbor atoms. Noble metals are represented in Nb 3 -NM compounds. These phases are superconducting with relatively high critical temperatures increasing in the order osmium, iridium, rhodium, platinum, and gold (Cahn and Haasen, 1983). In the Ni-As structure (B 8 structure), noble metals are combined in equiatomic proportion with transition metals. These compounds show a striking range of homogeneity due to the presence of three different interstitial sites (octahedral, tetrahedral, and trigonal), which can be occupied by the more electronegative metal components in different filling grades. In corresponding noble metal compounds, the transition metal atoms occupy a closepacked hexagonal array. The noble metals fill the octahedral voids. They form noble metal layers, which are placed between layers of the transition metal atoms in a simple hexagonal array with six nearest neighbors. While for the pure NiAs compound the axial ratio c/a is ^1.6, noble metal compounds of that structure type have axial ratios in the range of 1.39 to 1.28 (Table 9-28). The smaller distance between

the noble metal atoms contributes to an increased metallic character. Silver does not form phases of the NiAs type because of its more tightly bound d-shell compared to copper and gold. The NiAs structure can be transformed by small movements of atoms into a similar orthorhombic B31 structure (HumeRothery, 1956, 5). It is represented in noble metal compounds with 3B-, 4B-, and SBelements. In compounds of the CaF 2 structure (Cl structure), noble metals can occupy either the A- or the B-position. In the Aposition they occupy an f.c.c. sublattice. The structure may be regarded as a kind of electron compound with an e/a ratio of 8/3. CaF 2 structures with noble metals in the B-position ("anti-isomorph") are formed with SB-elements as compounds with a pronounced nonmetallic character. The behavior of the PGMs is rather complex both in terms of solid solubilities and intermetallic phase formation. Ruthenium, iridium, palladium, and platinum form very stable compounds with the refractory transition elements, especially zirconium and hafnium (Brewer, 1967; Darling, 1967, 1969). Platinum takes up to 18

Table 9-28. Axial ratios in Ni-As phases (Hume-Rothery, 1956, 9).

s Se Te As Sb Bi Ge Sn Pb Ga In

Ti

V

Cr

1.75 1.75 1.66

1.73 1.67 1.60

1.67 1.63 1.52 1.56 1.33

Mn

Fe

Co

Ni

1.63

1.68 1.64 1.49

1.54 1.46 1.38 1.43

1.25

1.34

1.55 1.46 1.35 1.37 1.39 1.31 1.31 1.27 1.25

1.53 1.40 1.42 1.24

1.25 1.23

1.27 1.23

Cu

Rh

Pd

Pt

Au

1.37

1.32

1.37 1.39 1.21

1.28

1.29 •

1.28 1.25 1.23

Ir

1.23

1.40 1.39

1.32 1.28

1.28

504

9 Noble Metals and Their Alloys

at.% zirconium into solid solution and even reacts with ZrC and ZrO 2 to form the hexagonal compound Pt 3 Zr. Relatively high solid solubilities exist for the PGMs in molybdenum and tungsten, which have been found in satisfactory agreement with theoretical calculations based on correlations between the (s + p) electron concentrations of the b.c.c., h.c.p., and f.c.c. phases (Brewer, 1963; Hume-Rothery, 1967). The PGMs have a high affinity for metals of the rare-earth group. They form numerous intermetallic compounds of different compositions and different structures. In view of the exceptionally high thermal neutron absorption cross sections of samarium, europium, and gadolinium, these PGM-rare-earth metal alloys are of special interest for the atomic industry (Funston, 1961; Loebich, 1971). 9.3.2.3 Superlattices Superlattices formed by disorder-order transformations are found in many noble metal alloys. Their generation is due to a marked difference in the alloying components, either in their electrochemical properties or in their atomic sizes. Stable superlattices are therefore not formed in systems like Ag-Au or Ag-Zn, but in systems like Au-Zn and Ag-Mg. The most frequently formed structures are of the compositions AB or AB 3 , which derive from the f.c.c. lattices in the disordered state. Depending on the degree of deviation between the atomic diameters, the rearrangement effects changes of the lattice dimensions, i.e., towards tetrahedral or rhombohedral structures (Table 9-29), and can be accompanied by considerable changes of the physical, mechanical, and chemical properties. Table 9-30 gives some typical examples of superlattice-forming noble metal

systems and the corresponding superlattice structures (Hirabayashi, 1991). The ordered structures always have higher electrical conductivity than the corresponding disordered states (Fig. 9-11; Savitskii and Prince, 1989). Similar behavior is shown by the thermal expansion. The highest increases in the hardness and magnetic properties are found at intermediate states, where maximum strains exist due to the combination of two different lattice structures. Ordering processes in such systems can be performed by quenching from the disordered state (which usually brings the alloys into an easy to deform state), followed by proper annealing at a suitable temperature and time to the desired state. Ordering processes with changing lattice symmetry are first order transitions with similar characteristics to those found in magnetic transition processes. Order-disorder transformations in noble metal alloys are of practical importance for the production of hard alloys (gold-copper, gold-silver-copper, platinum-copper, palladium-copper) and of special interest for strongly magnetic platinum-cobalt and platinum-iron alloys. In the system gold-copper, superstructures are formed at compositions Cu 3 Au, CuAu, and AuCu 3 (Fig. 9-12). 18 _14 E

j*10

Au 10

30

50 Cu (at.*?

70

90 Cu

Figure 9-11. Electrical resistivity in the Au-Cu system: 1) as-quenched, 2) air cooled (Savitzkii, 1989).

505

9.3 Alloys

Table 9-29. Changes in structure by superlattice formation. Fundamental structure

Composition

Superlattice structure Ll 0 Lli B2 B19

1:1 Al C\5-Fm3m

LI 2 (cubic) DO 22 (tetragonal)

3:1/1:3

A2

(tetragonal) (rhombohedral) (cubic) (orthorhombic)

L2j B2

1:1

(cubic) (cubic)

Djh-C4/mmm D^d-R3m 0;-Pm3m D^-Pmcm Oi-Pm3m

Oj-Pm3m

Table 9-30. Compositions and structures of superlattices in NM alloy systems (Benner, 1991, 1). Atomic ratio 3:1 or 1:3

Composition Pd 3 Fe Pt3Fe Pt3Ti Rh 3 Mo Rh3W Ir 3 Mo Ir 2 W Pd3V Pd 3 Nb Pt3V

2:1

1:1

Pd2V Pt2V Pt 2 Mo PdFe PtFe PtV PtNi PtCo PtCu PdCu

RhMo IrMo PtMo IrW PtNb

PdCu3 PtCu 3 PtNi 3 PtMn 3

Au3Cu AuCu3 Au3Pd Ag3Pt

Superlattice structure

Fundamental structure

Ll 2

f.c.c.

Ag3ln D0 19 (Au2Mn)

D0 22

f.c.c.

Au3Cd Au3Zn Ag3Mg

D0 23

f.c.c.

Ni2Cr

f.c.c.

Ll 0

f.c.c.

AuCul

AuMn AgCd AgZn

LI, L2 0 or B2

f.c.c. b.c.c.

h.c.p. B19

506

9 Noble Metals and Their Alloys 1100 ••-'-•

1084.87°C

1064.43°C 1000

0 Au

(a)

20

30 40 50 60 Weight percent copper

70

80

90

100 Cu

structure in the solid solution Al range

superlattice structures

Cu3AuI(cubic Ll 2 ) Cu3Au]I(cubic L l2)

(b)

APS

CuAu Korthorhombic LL) CuAul(tetragonal LL) (APB,M=5)

CuAu3 (cubic Ll 2 )

APB

• Cu O Au

(c)

Structure of anti-phase - boundaries in CuAul

Figure 9-12. Compositions and structures of superlattice phases in the Au-Cu alloy system: (a) Phase diagram of the system gold-copper (Massalski, 1983). (b) Crystal structures and atomic arrangement, (c) Structure of the anti-phase boundary of CuAuII.

9.3 Alloys

The ordered phases Cu 3 Au and CuAu form, over special temperature ranges, long-period "antiphase" superstructures which derive from the L l 0 and Ll 2 structures by an alternating, regular, linear arrangement of each of the nine or five single cell domains respectively in different orientations (Tachiki and Teramoto, 1966). The antiphase boundaries (APB meaning out of phase in the order row) are formed during the growth of ordered domains when they meet one another. This can affect the lattice geometry. The number of cells per, period depends on the valence-electron concentration (e/a) and is changed by alloying with different elements and compositions (Sato and Toth, 1965). High resolution electron microscopy has, in addition to X-ray methods, brought detailed information on the size and distribution of the microdomains (Hiraga, 1980). In gold-manganese alloys, changing the manganese content leads to a sequence of different structures from Au 4 Mn (Dlatype) -» Au 22 Mn 6 -• Au 31 Mn 9 (-> Id-antiphase structure based on the DO22-type) (Terasaki et al. 1981). Similar structures are also found with changing compositions in the system Pt-Ti, in the composition ranges Pt8Ti -> Pt 13 Ti -» Pt 21 Ti -> Pt 3 Ti (Schryvers et al. 1983). In the composition range 10-25 wt.% Pd, palladium and copper form the tetragonal L l 0 structure, and in the range 30-50 wt.% Pd the B 2 structure. If quenched, the alloys show excellent cold workability. In the Cu-Pt system, a long-range-ordered rhombohedral lattice is formed below 812 °C. The equiatomic Pt-Co alloy is of great importance for permanent magnets. Ordering occurs in such a manner that cobalt and platinum atoms arrange on alternate planes. Considerable stresses are caused during the ordering process by the coherence of tetragonal platelets on the

507

{110} planes of the cubic lattice. This can be seen from the Widmannstatten structure aligned with the {100} planes of the disordered lattice. The best magnetic characteristics are developed if the ordered and disordered phases are present in comparable proportions. Pt-Cr alloys show a continuous change from the Cu 3 Au type to the AuCu type in the range from 17-65 wt.% Cr. The phase Pt-Cr ranges from 43-84 wt.% Cr and has its magnetic saturation at ^30 at.% Cr. In the Pt-Mo system, at a composition with 30 at.% Pt, a superconducting ordered phase of the DO 1 9 type is formed. Here, also, the partially ordered state has the maximum effect with the highest critical transition temperature (7.5 K) (Fliikiger et al. 1972). Ordered A 3 B phases of Pd-Mn and Pd-Fe alloys show considerably higher solubility for hydrogen than the corresponding disordered phases. In Pd-Mn alloys, hydrogen helps the ordering process to proceed at lower temperatures (Flanagan and Sakamoto, 1993).

9.3.3 Selected Phase Diagrams The phase diagrams of noble metal alloy systems show a variety of different relations for alloys of the noble metals with each other, and systems with A- and Bgroup metals. Detailed presentations of most binary and numerous ternary noble metal alloy systems are given in many standard publications. Figure 9-13 shows some typical examples of phase diagrams of binary alloy systems forming continuous solid solutions, ordered phases, miscibility gaps, eutectic systems, and complex phases. In systems forming continuous solid solutions, three characteristic types, determined by the difference in the atomic diameters of the alloy components, can be

10

20

30

40

50

60

70

Weight Percent Germanium

Figure 9-13. Phase diagrams of binary noble metal systems: a) continuous solid solution, b) superlattice formation, c) eutectic systems, and d) complex systems.

20

30

40

Weight Percent Rhodium y' 1600-i

1400

-

800^

600-

y'''' 400^

200-

i'i

000 :

800-

600-;

11—TV 400-

200

400-

il.

.1

800^

o-

0

Weight Percent Tin

60

90

1

1

V

\

1

\

\

510

Noble Metals and Their Alloys

A

(0

(b)

(a)

%B

B

A

%B

B

A

%B

B

Figure 9-14. Types of systems for continuous solid solutions: (a) type A, (b) type B, and (c) type C.

distinguished (see Fig. 9-14 and Table 931). Silver-copper, with a difference in atomic diameters A & 12.5%, has a restricted solid solution and forms a eutectic-type system. In alloy systems with B-metals, the solid solubility in the a-range depends on the valency of the solute. Gold forms eutectic alloys with silicon, germanium, and tin with remarkably low melting points at the eutectic composition. The eutectic compositions of the corresponding silver alloys with silicon and germanium are of very similar composition (in at.%), but have slightly higher melting temperatures (Table 9-32).

Table 9-31. Systems forming continuous solid solutions. System

Difference in atomic diameter A (%)

Type of system

Au-Ag Pd-Pt Pd-Rh Pt-Cu Pd-Cu Pt-Ni Pd-Ni Au-Cu

0.2 0.2 2.2 6.9 7.2 9.2 9.4 11.4

A A A A + superlattice B + superlattice B + superlattice C C 4- superlattice

Table 9-32. Eutectic compositions and melting temperatures of copper, silver, and gold with silicon, germanium, and tin.

9.3.4 Influence of Alloying Elements

Si at.%

9.3.4.1 Basic Considerations Alloying with either noble metals or base metals can result in different degrees of dispersion: as solute atoms, as ordered phases, or as precipitates. Special properties are obtained by dispersion hardening or by forming composite materials. The extended capabilities obtained by alloying noble metals are used to influence, change, or adapt the mechanical, thermal, electrical, optical, and chemical properties over wide ranges according to requirements. Figure 9-15 shows some examples of the effect of alloy additions on hardening, and

Ge

Cu Ag Au

30 15.4 17.5

Sn

at.% Tm(°C) at.% Tm(°Q 802 830 370

36 25.9 27

640 651 356

7.7 11.5 20

798 724 280

Fig. 9-16 examples of effects of alloy additions on the electrical conductivity. 9.3.4.2 Metals Solid-Solution Hardening Table 9-33 gives some effective hardening elements for noble metals. In solid so-

9.3 Alloys

511

(a)

^80

"o 40

20

4

8

12 16 Alloy addition (%)

4O

20

o

Pd

4 8 12 Alloying element (%)

16

Table 9-33. Hardening elements for NMs.

5 10 15 20 Alloying element (%)

0 Au

Ru

Ir

0 Pt

2

4 6 8 10 12 Alloying element % )

14

16

Figure 9-15. Solid-solution hardening of noble metals: (a) silver (Butts and Coxe, 1967), (b) gold (Wise, 1964), (c) platinum (Zysk, 1979), and (d) palladium (Wise and Vines, 1979).

Noble metal

Hardening elements

Ag Au Pt Pd

Cu, Mn, Sn, Sb, AI, Mg, Pd Cu, Zr, Sn, Ge, Co, Ni, V, Be Ni, Rh, Ir, Cu, Fe, Ru, Os Ru, Os, Ir, Rh, Au, Fe, Co, Ni

lutions, due to elastic interactions between dislocations and solute atoms, higher flow stresses are required than in pure metals (Fleischer, 1964; Labusch, 1970,1981). Silver and gold alloys show characteristic solid-solution hardening behavior. The yield stress depends on the type and concentration of the solutes. They decrease with increasing temperature to a final value ("plateau stress", Fig. 9-17; Haasen, 1979, 1983, 1984). This is supposed to be caused by interactions between the dislocations and the solute atom groups, which cannot be released by thermal activation. The critical resolved shear stress, TC, of the solid solutions is characteristic for each type of solute. Their dependence on solute concentration c can, according to the theoretical approach of Labusch, be satisfactorily described by (see Fig. 9-18) Tc = A.

(9-13)

9 Noble Metals and Their Alloys

where A is a constant of proportionality. The different slopes depend on the type of foreign atoms. Precipitation Hardening

"

0

2 4 6 8 Alloying addition (at.%)

Systems with miscibility gaps and with decreasing solubility as a function of the temperature can be precipitation hardened. This can be done in duplex phase regions using miscibility gaps or the temperature dependence of the solubilities in the solid state. This is of practical importance for the hardening of construction parts of Ag-Cu alloys, and for the hardening of Au-Cu and Au-Ag-Cu alloys and platinum group metal alloys for different applications like jewelry, dentistry, or industrial apparatus parts and construction components. In binary and ternary gold alloy systems with higher copper contents, additional hardening can be obtained by ordering reactions (Yasuda and Hisatune, 1993).

10

(b)

0

2 4 6 8 Alloying addition (at.%)

(c)

9.3.4.3 Carbon in Plantinum Group Metals Cu

50

Ru 1=40
20 Pt 10

Rh

Alloying addition

Ag

The platinum group metals dissolve large quantities of carbon in the molten state, but its solubility in the solid state is very low. After solidification of liquid PGMs with a high carbon content, characteristic microstructures are formed. Pt-C alloys show lamellar graphite flakes (Fig. 9-19) resembling those of cast iron, causing brittleness and cracks when deformed. In palladium, carbon is precipitated in a spherical form (Fig. 9-20). This is the reason that Pd-C alloys, though much harder than Pt-C alloys, are relatively ductile compared to Pt-C alloys (Darling, 1973, 2). Figure 9-16. Influence of alloying elements on the electrical conductivity: (a) silver, (b) gold, (c) platinum, and (d) palladium (Doduco, 1977, 1-4 resp.).

9.3 Alloys

100

200 300 400 Temperature T (K)

500

Figure 9-17. Critical shear stress of silver-aluminium alloy single crystals (Haasen, 1984, 3).

513

All PGM-carbon alloys have slightly lower melting temperatures than those of the pure metals. The melting temperatures decrease in each period with increasing atomic weight (Table 9-34; Nadler and Kempter, 1960). The solubility of carbon in solid platinum is very low (<0.02% by weight at 1700°C), but it diffuses rapidly through platinum even at 1200 °C. Palladium takes 0.04 wt.% C at 800 °C and ^0.45% C at 1400 °C into solid solution, this is accompanied by a slight increase in the lattice parameter from 3.894 A to 3.920 A (0.3894-0.3920 nm). Carbon dissolves interstitially and effects a considerable increase in the hardness. 9.3.4.4 Grain Refining Coarse-grained structures are formed by excessive crystal growth during solidifica-

Au-Ga

Table 9-34. Melting temperature of carbon-containing PGMs (Darling, 1973, 3). Metal

0.20

Figure 9-18. Critical shear stress of solid-solution hardened gold alloys (Haasen, 1984, 4).

Ru Rh Pd Os Ir Pt

Melting point

Carbon eutectic point

2310 1963 1554 3033 2447 1772

1942 + 16 1694+17 1504+16 2732 + 22 2296+16 1736+13

0.86 0.88 0.97 0.90 0.95 0.98

Figure 9-19. Pt-C alloy rnicrostructure; x 45 (Darling, 1970). Figure 9-20. Pd-C alloy rnicrostructure; x 45 (Darling, 1970).

514

9 Noble Metals and Their Alloys

tion from the melt and on recrystallization during annealing treatments of workhardened material. Due to the fact that plastic flow in crystals is orientation-dependent, further deformation like deep drawing or bending can cause the formation of undesirable, rough surface structures ("orange peel effect"). To avoid this, fine-grained materials are needed. Finegrained noble metal alloys have higher mechanical strengths and advantageous working behavior, and show a much better surface appearance after plastic deformation processes, also after etching, polishing, and galvanic coating. Grain refining and grain stabilization are especially important for jewelry and dental alloys, but are also required for silver and platinum metal alloys used for construction components of chemical equipment, spinnerets, and seal rings. The basic reaction during the production of fine-grained cast material is the formation of finely dispersed, insoluble particles, which serve as nuclei for crystallization. They can be formed by the reaction of highly active elements with the crucible material (e.g., Zr, ZrC) or by reaction with alloy components. Grain growth during annealing can be suppressed by segregation of either the insoluble atoms or the solute atoms with a large dimensional misfit into the grain boundaries (Riabkina and Gal-Or, 1984). Effective grain-refining additions for castings of gold and gold alloys (Au-Cu and Au-Ag-Cu in 14 and 18 carat) are small amounts of iridium, ruthenium, rhenium, zirconium, barium, tantalum, and niobium, (0.01-0.1 wt.%), and combinations of Co/Ba, Co/Mo, and Ni/Mo. Zirconium, barium, and cobalt increase the recrystallization temperature by 150200 °C and suppress grain growth on further annealing. In 18 Ct Au-Ag-Cu alloys,

iridium and rhenium must be added prealloyed with palladium. Tin (0.3 wt.%) and antimony (0.1 wt.%) decrease the recrystallization temperature of 18 Ct goldcopper alloys by «100 °C. This is thought to be caused by possible influences of these elements on the ordering processes during the formation of superlattice structures (Ott and Raub, 1980/81/82). The most common grain-refining additive for silver is nickel, alloyed at a concentration of 0.15 wt.% ("fine-grain" silver). Gold, zinc, aluminum, antimony, and cadmium increase the recrystallization temperature. Grain refining of the platinum metals is usually performed by the dispersion of insoluble inclusions which also act as agents for dispersion hardening. Alloying with 1 wt.% niobium, titanium, or zirconium increases the recrystallization temperature to 1000 °C, 1200 °C, and 1300°C, respectively. This grain-refining effect is ascribed to the formation of finely dispersed refractory phases by reaction of the additives with traces of oxygen (Darling, 1973, 3). 9.3.5 Noble Metals and Gases 9.3.5.1 Oxygen The noble metals, except gold, form solid and gaseous oxides. Oxygen dissolved in the noble metals has a strong effect on the hardness and working behavior. Silver, platinum, and platinum-group metal oxides are of practical importance in catalytic reactions. Very thin layers of Ag2O are formed on the surface of silver in oxygen or air below 180°C. Above this temperature the oxygen partial pressure of the Ag2O exceeds 760 Torr (0.10 M n m " 2 ) and the oxide decomposes. Even at temperatures up to 800 °C, no visible attack can be seen, though small weight losses can be measured. Oxygen dissolves in the silver

515

9.3 Alloys

lattice occupying interstitial lattice sites. Its solubility increases with increasing temperature; this increase is considerable at the melting temperature (by a factor of « 80), and then decreases with further increases in the temperature (Table 9-35). The solution of oxygen in liquid silver is favored by the formation of Ag 2 O, which dissolves in the liquid silver. The melting temperature of silver decreases with increasing oxygen pressure of the surrounding gas atmosphere according to (Raub, 1940, 1): 71 = 961.5-22.31

(9-14)

Gold shows, in contrast to silver, a slow increase in the melting temperature in an oxygen atmosphere, of 6°C for each kilobar (100 MNm~ 2 ) increase in the gas pressure. The platinum group metals form a variety of solid and gaseous oxides in different oxidation states. The solid oxides have tetragonal and hexagonal crystal structures (Table 9-36; Darling 1973, 4). Platinum forms three solid oxide compounds which are stable at different temperatures and exhibit different oxygen partial pressures. Solid PtO 2 in the hexagonal a-modification is used as a catalyst. The orthorhombic (3-modification is insoluble in all acids including HNO 3 /HC1 (aqua regia). The solubility of oxygen in platinum is very low. Therefore even very thin coatings of platinum on reactive metals are an effective protection against oxidation. Accelerated diffusion of oxygen through platinum electrodes has, however, been found under electrolytic conditions (Hoare, 1969). Platinum oxides are also stable in the gaseous phase at vapor pressures in the range of 10~8 to 10" 4 atm (10~ 3 to ION m~ 2 ) at temperatures above ^ 4 0 0 K (Fig. 9-21). The formation of volatile oxides is responsible for weight losses when

Table 9-35. Oxygen uptake by silver (DEGUSSA, 1967). T(°C)

200 400 600

800

973 1000 1100 1200

Solubility 0.03 1.4 10.6 38.1 3050 3000 2700 2500 (ppm)

Table 9-36. Structure of PGM oxides (Darling, 1973, 4). Oxide

Structure type

a-PtO2 P-PtO2

primitive hexagonal primitive orthorhombic primitive cubic tetragonal

3.08 4.486

PdRhO

hexagonal

5.22

6.0

Rh 2 O 3 (LT)a Rh 2 O 2

hexagonal (corundum) tetragonal (rutile)

5.108

13.87

Pt 3 O 4 PdO

a

Unit cell dimensions

4.19 4.537 3.138

5.585 3.043

4.4862 3.0884

LT = Low temperature modification.

Ru0 3

400

800 1200 1600 2000 Temperature (K)

Figure 9-21. Vapor pressures of PGM oxides (Darling, 1973, 8).

516

9 Noble Metals and Their Alloys

platinum metals or platinum metal alloys are heated in oxygen-containing atmospheres, causing remarkable weight losses in long term runs, e.g., in the catalytic oxidation process of NH 3 to NO. The vapor pressures of platinum and rhodium oxide are almost coincident, so that the composition of the catalyst gauze is not seriously changed by the formation of volatile oxides. The volatile oxides are recovered by leading the outcoming gas stream through a second gauze consisting of a goldpalladium alloy, on which the platinum and rhodium oxides are deposited. Rhodium and palladium form stable oxide skins when heated in air. Both also take oxygen into solid solution. The solubility increases in the order Rh -» Pd -> Ag. PdO has a tetragonal crystal structure and is semiconducting, RhO 2 is an electrical conductor. The electrical resisitivity decreases with decreasing temperature from l x l O ~ 4 f i c m a t room temperature, falling to 1 x 10" 5 |Q cm at 4.2 K. Fine powders of ruthenium and osmium oxidize readily at 20 °C to RuO 2 and OsO 4 , respectively. The latter is a pale yellow crystalline solid that melts at 39.5 °C. Above its melting point it exhibits a high vapor pressure, increasing from 25 Torr (3.33 kN m" 2 ) at ^40°C to 760 Torr (101 k N m " 2 ) at ^130°C. It is extremely poisonous. RuO 2 is a semiconductor. It is the basis of conductive and resistive inks for electrical applications. The solubility of oxygen in solid ruthenium is very low. Dissolved oxygen in the melt is rejected on solidification as solid oxide, which is located on the grain boundaries. Ruthenium sponge usually has an oxygen content between 2000 and 4000 ppm. It can be reduced by sintering in hydrogen to ^ 7 0 - 5 0 ppm. Further reduction of the oxygen down to 10-3 ppm is possible by electron-beam melting (Darling, 1973, 5).

9.3.5.2 Hydrogen All noble metals, except palladium, take only very small amounts of hydrogen in solution. Platinum dissolves 15 ppm hydrogen at 900 °C and 760 Torr (101 kN m- 2 ). Palladium and certain palladium alloys are unique in their capability to dissolve large amounts of hydrogen of up to « 2800 times their own volume (Moore, 1939; Lewis, 1967,1994). This effect finds practical use for hydrogen permeation membranes in hydrogen purification equipment. Hydrogen enters palladium interstitially. The hydrogen atoms are located in the octahedral holes of the f.c.c. lattice. Above 295 °C, palladium and hydrogen form a continuous series of solid solutions with an f.c.c. structure. Below this temperature, the phase splits into an f.c.c. palladium-rich a-phase and an f.c.c. hydrogen-rich (3-phase, forming a miscibility gap which broadens with decreasing temperature. The equilibrium hydrogen presure at 295 °C amounts to 19.87 atm (2.0 MN m" 2 ); the composition at the critical point corresponds to palladium with 21 at.% hydrogen. The P-phase takes hydrogen up to 1300 times the volume of the palladium, corresponding to a composition of palladium with 50 at.% hydrogen. Further quantities of hydrogen, up to a total of 2800 times the volume of the palladium, can be loaded by cathodic deposition. The two-phase region has, according to the phase rule, constant vapor pressures at constant temperatures, which amounts to 1 atmosphere (0.1 MN m~ 2) at ^140°C. The occupation of interstitial sites by hydrogen atoms causes expansion of the palladium lattice. The lattice parameters increase with increasing hydrogen content from 3.891 A (0.3891 nm) for pure palladium up to 4.06 A (0.406 nm) for an

9.3 Alloys

alloy containing 75 at.% H (Nelin, 1971). Thermal cycling of Pd-H alloys in the duplex phase region causes the development of high mechanical stresses due to different changes of the lattice dimensions of the two phases due to the temperature-dependent, changing amounts of dissolved hydrogen. Pure palladium, if repeatedly heated and cooled in a hydrogen-bearing atmosphere therefore becomes brittle and cracks. Palladium-silver alloys with 20-25 wt.% silver dissolve higher amounts of hydrogen than pure palladium without becoming brittle, are more resistant to thermal cycling, and are also more permeable to hydrogen at lower temperatures (Lewis, 1982). Hydrogen purification equipment based on the alloy palladium with 23 wt.% silver is widely used in the semiconductor industry to produce ultrapure hydrogen of up to 99.99999% purity. Palladium-silver permeation membranes are provided in the form of bundled tubes with a wall thickness of ^75 jim, sealed at one end. Compressed hydrogen is passed from outside the tubes through the walls of the membranes, dissociating and diffusing through the alloy to the inner side of the tubes, leaving most impurities outside the membrane (Hunter, 1966; Knapton, 1977). The solubilities of the hydrogen isotopes in palladium and palladium-silver alloys differ to an appreciable degree. Diffusion membranes of palladium or palladium-silver alloys can therefore be used for separation and concentration processes of different hydrogen isotopes. The equilibrium hydrogen/deuterium separation factors vary from ^2.4 at 25 °C to 3.7 at - 8 0 ° C (Wicke and Nernst, 1964). For the hydrogen/tritium ratio a value of 2.8 has been measured at 25 °C (Sicking, 1970).

517

9.3.6 Liquid-State Properties 9.3.6.1 General Knowledge of the mechanical and thermal properties of liquid metals is of importance for the understanding and control of the processing conditions on casting, e.g., the flow and wetting behavior and the volume changes on solidification. Knowledge of the electrical properties is essential to control the processes of melt electrolysis. Important properties in this respect are density, viscosity, surface tension, and electrical resistance. Measurements have been made on liquid noble metals and liquid noble metal alloys by X-ray and neutron scattering. In the solid state, metal atoms oscillate around fixed lattice positions. Metal crystals begin to melt if by thermal action the amplitudes of the thermal oscillations of the atoms exceed about 10% of the distance to their next neighbors. Above the melting point, the translational periodicity of the lattice is lost, but some short-rangeorder remains, limited to a few atomic diameters (Fig. 9-22). Strong atomic bonds, as exist in intermetallic compounds for instance, may still partially be present in the liquid state just above the melting temperature. The atomic volume increases on melting by 3 - 5 % ; the entropy of melting amounts to about 9 J mol~ 1 ; and the electrical resistance increases by a factor of 1.5-2. The distances of the atoms in the short-range-order arrangement change linearly with the composition at temperatures near to the solidification temperature, and are retained in the liquid state (Fig. 9-23; Richardson, 1974). 9.3.6.2 Density The density of liquid noble metals decreases with increasing temperature. The

518

9 Noble Metals and Their Alloys

Table 9-37. Density of liquid NMs (Poniatowsky, 1995).

10

Figure 9-22. Radial distribution of relative mass-densities in liquid gold: Vertical lines mark the radii of the coordination spheres in the lattice; their length corresponds the occupation density N(r) in this region (Schulze, 1974, Poniatowsky, 1995).

Metal Temp. Q 3 (°Q (g/cm )

Alloy

Ru Rh Pd Ag Os Ir Pt Au

Ag-Cu40 Ag-Bi40 Ag-Pd40 Ag-Sn40 Au-Cu40 Au-Fe40 Au-Ge40 Au-Sn40 Pd-Co40 Pd-Cu40 Pd-Ni40 Pd-Rh40 Pt-Co Pt-Cu40 Pt-Fe40 Pt-Ni40 Rh-Fe40 Rh-Co40 Rh-Ni40 Rh-Sn40

2250 1960 1552 961 3027 2443 1770 1064

10.9 10.70 10.49 9.35 20.1 19.39 18.91 17.19

Temp. Q 3 (°Q (g/cm ) 1100 1055 1600 1000 1100 1550 1200 1000 1600 1600 1600 2000 1800 1800 1800 1800 2000 2000 2000 2000

8.78 9.53 9.60 8.11 14.40 13.27 11.80 12.06 9.27 9.24 9.30 10.0 12.80 15.0 14.25 14.87 9.15 9.23 9.28 8.99

influence of alloying elements can be estimated for simple systems by the rule of mixtures. Selected values of the densities of liquid noble metals and NM alloys are given in Table 9-37. 9.3.6.3 Viscosity

(e)

(f) 0.2

0.3

0.4

0.5

Sin 0 / X

Figure 9-23. X-ray intensities of liquid Au-Sn alloys: (a) Au at 1393 K, (b) Au75-Sn25 (at.%) at 699 K, (c) Au70.6-Sn29.4 (at.%) at 568 K (eutectic), (d) Au60Sn40 (at.%) 685 K, (e) Au50-Sn50 (at.%) at 698 K (V), and (f) Sn at 505 K; 6 = angle of incidence, X = wavelength. Short-range order is indicated by interatomic distances in the composition range above 25% Sn (Richardson, 1974, Poniatowsky, 1995).

The viscosity of liquid noble metals is material-specific and temperature-dependent. Influencing factors are the interatomic bonding forces, atomic sizes, and coordination numbers in short-range-ordered clusters. The viscosity of liquid alloys also depends on the type and concentration of the alloy components. Selected values of the viscosity of liquid noble metals and some alloys are listed in Table 9-38.

519

9.3 Alloys

Table 9-38. Viscosity of liquid NMs and NM alloys (Poniatowsky, 1995). Elements Metal

NM alloys

n

Temperature (°Q

(10~3 Pas)

Ag

1000

3.94

Au

1100

4.56

Pd

1560

4.22

Pt

1797

7.60

Rh

2000

5.11

9.3.6.4 Surface Tension

The surface tension of liquid noble metals in inert gas atmospheres is in the range 350-2500 mJ m~ 2 (compare, with H 2 O at 20°C at 70 mJ m" 2 ). It is also temperature-dependent and influenced by alloying elements. The surface tension is primarily of interest for soldering operations in relation to the flow and wetting behavior. A relation has been established with the heat of evaporation regarding the work needed to separate or remove surface atoms (Takeuchi, 1972). Selected values of the surface tension of some noble metals and some noble metal alloys are listed in Table 9-39. The wetting behavior is determined by the surface tension of the liquid metals, and may be influenced by surface reactions with the support to be wetted. The spreading speed of the diameter of drops of liquid silver and gold on pure metal surfaces increases linearly with time. The spreading speed is temperature-dependent and follows an Arrhenius relation with activation energies comparable to those of diffusion.

Composition

Temperature (°C)

n

(10"3 Pas)

Ag-Cu40 Ag-Au40 Ag-Ge40 Ag-Sn40

1000 1100 1200

Au-Cu40 Au-Sn40

1100 900

4.56 2.15

Pd-Ni40 Pd-Cu41.8 Pd-Su40

1530 1540 1450

5.07 6.66 3.83

Pt-Co40 Pt-Ni40 Pt-Cu40

1800 1800 1800

4.82 4.80 5.25

3.72 4.17 1.50 2.02

900

Table 9-39. Surface tension of liquid NMs: (a) NMs and (b) NM alloys (Poniatowsky, 1995). (a)

7TC)

Metal

Q

961 1064 1552 1770 1960

Ag Au Pd Pt Rh

Surface tension, [mJ • m 2]atr(°C)

7TQ

r+200(°Q

919 1138 1457 1746 1915

898 1100 1419 1715 1849

Composition T Q Composition T Q Ag(at.%) (°C) (mJ-m 2 ) Au (at.%) (°C) (mJ-m 2 ) Ag-Cu20 Ag-Au20 Ag-Pb20 Ag-PdlO Ag-SnlO

1000 1000 1000 1600 1000

860 967 560 900 780

Au-Co20 Au-NilO Au-Cu20 Au-Si20 Au-Sn20

1500 1500 1200 1300 1150

1121 1100 1120 965 845

Composition T Q Composition T Q Pd(at.%) (°C) (mJ-m 2 ) Pt (at.%) (°C) (mJ-m 2 ) Pd-ColO Pd-NilO Pd-CulO Pd-PtlO

1600 1600 1600 1800

1465 1460 1440 1345

Pt-ColO Pt-Ni20 Pt-Cu20 Pt-RhlO

1800 1800 1800 1966

1720 1700 1605 1717

520

9 Noble Metals and Their Alloys

Measured values are between 32 kcal (134 J) (for copper on platinum) and 83 kcal (347 J) (for gold on nickel) (Nicholas and Poole, 1966).

Table 9-40. Electrical resistivity of liquid NMs: (a) noble metals and (b) noble metal alloys (Poniatowsky, 1995). (a)
Metal

9.3.6.5 Electrical Resistance

atTTQ

The electrical resistance of pure liquid noble metals increases with increasing temperature and can be described in analogy to that of solid metals in the first approximation by

Ag Au Pd Pt

961 1064 1600 1770

at r+100(°C) 17.25 31.20 120 101.5

18.0 32.66 103.0

(b)

Qc = * T + p

(9-15)

where oc and b are coefficients. Selected values are listed in Table 9-40. The thermal conductivity {X in J cm" 1 s" 1 K~ 1 ) can be estimated using the Wiedemann-Franz law I = ae T c

(9-16)

where <7e is in Q"1 cm" 1 , T is in K, and

9.3.7 Special Alloys 9.3.7.1 General A great range of noble metal alloys has been developed, for a broad range of different applications. Detailed values of the mechanical, physical, and chemical properties are reported in numerous corresponding reference books (see General Reading). In the following a selection of alloys of major technical importance are presented. (All concentrations noted here are given in wt.%). 9.3.7.2 Silver Alloys The alloys Ag-Ni0.15 (fine-grain silver), Ag-Sil.5 (hard silver), and Ag-Cd5 exhibit excellent corrosion resistance and high electrical conductivity. They are therefore preferentially used as construction materials for equipment and apparatus in the chemical industry.

ee(10-8Qm)

Composition (at.%) Ag-Cu20 Ag-Au20 Ag-SilO Ag-GelO Ag-SnlO

1000 1135 1600 1000 970

19.4 39.3 32 70 60

Au-Ni20 Au-Cul5 Au-Ga20 Au-Sn20 Au-Bi20

1200 1800 1000 600 800

61.9 45.5 68.2 62 95.0

Pd-FelO Pd-ColO Pd-NilO Pd-Cu30 Pd-AglO

1600 1600 1600 1600 1600

120 127 120 88 100

Ag-Cu alloys: Copper effects hardening by three different mechanisms: solid solution, improved rate of work hardening, and, due to the miscibility gap, hardening by precipitation. The compositions of the majority of industrially important alloys are on the silver-rich side of the eutectic point. Ag-Cu alloys with 3-20 wt.% Cu (hard silver alloys) are used for contacts, switches, circuit breakers, and motor contactors. Ag-Cu7.5 (sterling silver), Ag-Cu4.2 (Britannia silver), and Ag-CulO (coin silver) are used for flat and hollow tableware, jewelry items, coinage, and electrical con-

521

9.3 Alloys

tacts. Ag-Cu28 (eutectic) and ternary or multicomponent alloys with Zn, Cd, Sn, Li, and Ni) find wide application in brazing alloys. Ag-Cu-P alloys are braze fillers for flux-free brazing. Ag-Mg0.25-Ni0.2 is used for electrical contacts that have to be affixed by brazing without loss of hardness, for high thermal conductivity spring clips for vacuum tubes, relays, springs, and similar parts. Ag-Mg-Ni alloys harden when heated in air by internal oxidation of the magnesium. The oxides are submicroscopically distributed in the matrix (dispersion hardened). This process is possible because the oxygen diffuses inwards faster than the magnesium diffuses outwards. Nickel prevents grain growth. The electrical conductivity is 75% IACSat20°C. Ag-Pd alloys of different compositions serve as electrical contacts; Pd-Ag 50-70 % is used as an electrodiode layer in capacitors and as a conductor material in electronic devices. These alloys are also used for brazing stainless steel, nickel-based alloys, and other heat-resistant alloys. Ag-Hg alloys play an important role in dentistry. Silver alloys with tin or lead as binary, ternary, or multicomponent alloys with indium, tin, copper, and zinc are used as soft solders for general purposes and in special compositions as solders for electronic devices. Ag-Ca alloys are the basis for the production of skeletal silver catalysts. Ag-Cu-Au alloys (see Sect. 9.5.2.1) are used in jewelry and dentistry in a large variety of compositions. Because of its high thermal-neutron capture cross-section and good corrosion resistance in hot water, Ag-Cd5-Inl 5 is used as the cladding material for control rods in nuclear fuel reactors.

9.3.7.3 Gold Alloys Au-Ag20 is used for low voltage electrical contacts, e.g., rivets. Au-Co5, Ag-Ni5, and Au-Ag26-Ni3 are used as contact materials which are resistant to silver migration. Au-Ag-Cu alloys (colored-gold alloys) are the most important basis for the selection of jewelry and dental alloys with defined colors. They contain, in general, additions for grain refining. Their melting temperatures are adjusted by additions of zinc or nickel for jewelry alloys, and palladium or platinum for dental alloys. The ternary system has a broad miscibility gap which extends from the Ag-Cu side to the gold corner. Figure 9-24 shows isothermal sections through the miscibility gap. The broadening of the miscibility gap with decreasing temperature enables precipitation hardening in these composition ranges (Fig. 9-25). In addition, strengthening occurs in the higher gold-composition range (>75 wt.% Au) by formation of the ordered Au-Cu phase. Au-Ni-Cu alloys (white-gold alloys) are jewelry alloys, and

\

Miscibility gap after solidification

Ag

20

40 Cu, (wt.%)

80

60 -

Figure 9-24. Isothermal sections through the miscibility gap in the Au-Ag-Cu system (Ullmanrfs, 1990).

9 Noble Metals and Their Alloys

0

10

20

Silver.%

30 0

10

20

30

Silverf%

40 0

10

20

30

40

50

60

Sitver,%

Figure 9-25. Variation of hardness with silver content in the Au-Ag-Cu system (Sistare and McDonald, 1979).

are also used over typical different composition ranges. In this system, higher hardness values are reached by annealing (Yamauchi et al. 1981). Au-Ag-Pd alloys are used for jewelry objects if nickel may cause allergic reactions. They have higher melting temperatures but are relatively soft and cannot be age hardened by precipitation. They may be made susceptible to precipitation hardening by additions of copper. Au-Pt alloys are, because of their high corrosion resistance, used for spinnerets in rayon production, and serve also as high melting platinum solder. A typical composition is Au-Pt30 with a liquidus temperature of 1450 °C and a solidus temperature of 1228 °C. The alloys with 40-65% Au can be effectively age hardened by quenching from 1100°C and annealing at ^500°C to tensile strengths at room temperature of the order of 1400 N/mm 2 . Segregation effects occurring with these alloys during ingot production are encountered by alloying with third elements (e.g., 0.5 wt.% of either rhodium, ruthenium, or

iridium). They broaden the miscibility gap and turn the system into the peritectic type by displacing the two-phase region up to the solidus. Au-PtlO and Au-Ag25-Pt5 are alloys for electrical contacts working under highly corrosive conditions. Au-Cul4Pt9-Ag4 is used for oxide-free slip contacts of high reliability at very low voltage and current, for instance as test value transmitters in the form of bent wire pieces or cut foils. Au-Ni-alloys form continuous series of solid solutions at high temperatures. These alloys are used as brazing alloys. The most common composition is Au-Nil8. Au-(Fe, Co, Ni) alloys with small amounts (up to 1-3%) of these solutes are used as hard (galvanically deposited) surface coatings on electrical contacts. 9.3.7.4 Platinum Alloys Platinum-palladium alloys are used in jewelry and with further additions (e.g., gold, rhodium) for electrical contacts and production components in glass manufac-

9.3 Alloys

turing and machine and apparatus parts that have to withstand strongly corrosive media. Platinum-iridium alloys find different applications in electrochemistry, chemical technology, medicine, and jewelry. Iridium contents (in wt.%) are 0.4-0.6 for laboratory ware, 5-15 for jewelry, 10 for medical applications, 10-25 for electrical contacts, 10 for electrodes, and 25-30 of pens, needles, and springs. Maximum hardness (^360 HB) is obtained with 30% iridium. Platinum-rhodium alloys are used with 10% rhodium as a catalyst in the oxidation of NH 3 , as the standard element in thermocouples (Pt-Rh of different compositions/Pt), and as heater elements in electrically heated high-temperature furnaces. Pt-Rh alloys with additions of gold are especially resistant to wetting by molten glass. In platinum-ruthenium alloys ruthenium strengthens platinum very effectively; with 10% ruthenium the hardness increases to 168 HV. The alloy Pt-Ru5 has the same strength, hardness, and corrosion resistance as the alloy Pt-IrlO (hard platinum). Applications include jewelry items, electrodes for electrochemistry, spark plug electrodes, electrical contacts, and wearresistant wires in potentiometers. Higher alloy compositions of up to 14% ruthenium and ternary types with 15 % rhodium and 6% ruthenium are used for electrical contacts and potentiometer wires for heavy duty applications. Platinum-nickel alloys have been used as electrical switches and material for high strength requirements and high temperature resistance. Practical compositions reach from trace amounts up to 20 % nickel with a maximum tensile strength of ^1725 MPa after 90% cold reduction. Platinum-tungsten alloys, commonly used in compositions with 5-8 wt.% Co,

523

have excellent wear resistance and are therefore used for potentiometers, glow wires in copiers, diode switches, heater elements, and claddings on rods for fuel elements in atomic reactors. Platinum-cobalt alloys with compositions near 50 at.% Co form strong permanent magnets with coercivities of up to 540 kA/m at room temperature. Typical uses are focusing magnets, hearing aids, miniature motors, medical instruments, and implants and films for digital magneto-optical recording. Platinum-gold and platinum-gold-rhodium alloys have much higher strengths than pure platinum at 1000 °C, a finer grain structure, and better resistance against wetting by molten glass. These alloys are therefore particularly suitable as materials for crucibles used in chemical analytical operations. 9.3.7.5 Palladium Alloys Pd-Ag23 exhibits extremely high hydrogen permeability and is used for semipermeable membranes for the extraction and purification of hydrogen. The Pd-Ag alloy with 40 at.% Ag has the highest electrical resistivity at room temperature. Because of its characteristic, flat, s-shaped curve for electrical resistivity versus temperature over an extended temperature range from 20 °C to 800 °C, it is widely used as a material for precision resistors in electrical technology. Its specific resistance between 20 °C and 100 °C is nearly constant and amounts to 42.5 \iQ cm; the specific resistance between ^20°C and 800 °C changes from 42.5 to 44 \xQ cm. Palladium-copper (15-40 wt.%) alloys serve as surface material for electrical contacts in the milliampere range, and as hard, wear resistant material for slip rings with high conductivity. Pd-Ag-Cu alloys con-

524

9 Noble Metals and Their Alloys

taining ^ 4 5 % Pd can be age hardened. These alloys find applications in dentistry. An addition of 1 % Zn accelerates age hardening. Ternary alloys containing 1025 % Pd are used as hard solders for the formation of high-strength brazed joints. Palladium-gold alloys are the basis of precision resistance alloys. An addition of iron accentuates the formation of long range ordering. The alloy Au50-Pd45Mo5 has a specific resistance of 100 |i Q cm with a temperature coefficient of resistance (TCR) of 0.00012/°C between 0 and 100 °C; the tensile strength can reach 1060 N/mm 2 . Palladium-ruthenium alloys find different applications, for instance, with 4.5% ruthenium as a standard jewelry alloy in the United States, with ruthenium contents of up to 12 % as a material for electrical contacts, such as catalysts for the reduction of nitric oxide. 9.3.7.6 Solders In soldering and brazing processes, materials are joined by applying liquid metals temporarily between the surface regions to be bonded. Process (i.e., working) temperatures are above the liquidus temperature of the solder, but always below the solidus temperature of the materials to be bonded. After cooling and solidification, stable bonds are formed. In order to wet and spread over the surface, the molten alloy must be able to form solid solutions or intermetallic compounds with the solid material. Low surface tension and good flow properties are further requirements to obtain void-free filling of the capillary space between the bonded faces. This is an essential precondition for the formation of mechanically stable bonds with optimum thermal conductivity. Soldering and brazing processes need clean and pure metallic

(oxide-free) surfaces. These are provided by applying flux materials in the solder operation or by working in protective or reducing gas atmospheres (N 2 , H 2 , mixtures of N 2 and H 2 , Ar) or in a vacuum. Solder materials are generally classified by their melting temperatures or their working temperatures as: • soft solders - liquidus temperatures below 450 °C; • hard solders - liquidus temperatures between 450 °C and 950 °C; and • high temperature solders - liquidus temperatures >950°C. Hard solder materials are also designed as brazing fillers. Noble metals containing soldering and brazing alloys are based on silver, gold, palladium, or platinum. They have high mechanical and thermal stability, high electrical conductivity, and corrosion resistance. They cover, over a large variety of alloy compositions, a wide range of different melting temperatures. These alloys are used to join noble metals (jewelry and dentistry), base metals, or ceramics with metals. Noble metal based solder alloys are usually designed in relation to the main constituent noble metal and the area of application (Krappitz et al., 1991). Soft Solder Alloys Most of these alloys are based on tin and lead. Additions of silver improve the mechanical strength and suppress leaching of silver surfaces during the soldering process. Silver-containing solders based on tin are applied in food canning, in the installation of copper tubes, and in cryotechnology. Silver-containing eutectic binary and ternary alloys of lead and tin form mechanically and thermally stable solder connections, and are therefore used as solders for the assembly of semiconductor devices.

9.3 Alloys

Table 9-41. NM-containing soft solder alloys. Type

Composition

Solidus Liquidus

Sn-based

Sn62Pb36Ag2 SnAg3.5 SnAg25SblO

179 E a 221 E a 228

Pb-based

Pb92.5Ag2.5Sn5 PbAg3 Pb92.5Ag2.5In5

280 E 305 E a 307 E a

Au-based

AuSn20 AuGel2 AuSi2

280 E a 356 E a 370

395

a

740

E = Eutectic composition.

Table 9-41 shows some typical soft solder alloy compositions (Beuers and Krappitz, 1995). Hard and High-Temperature Solder Alloys Solder alloys based on gold are important for jewelry, dentistry, and high temperature resistant solder connections. Caratage, color, mechanical stability, and corrosion resistance are most important for jewelry solders. The requirements for dental solders include their color, corrosion resistance, and mechanical strength. Brazing filler materials based on gold for industrial applications are Au-Ni and AuCu alloys. They have high oxidation resistance and good thermal stability up to 800 °C (Colbus and Zimmermann, 1974). The production of turbines, jet engines, and components and devices for nuclear and space technology are their fields of application. Silver-based hard solders are the biggest group of hard solder materials, and are preferentially used in industrial applications, electronics, and silverware. The melting ranges of these alloys extend from 610-960 °C. Additions of copper and zinc lower the melting temperature; cadmium

525

improves the strength, the flow and wetting behavior, and the ductility. Silvermanganese solder alloys are especially suitable for soldering hard metal and refractory metal surfaces (molybdenum, tungsten, tantalum). Supplementary additions of nickel or cobalt increase the strength of the bond. Ag-Pd-Cu alloys are used for high temperature applications (gas turbines, supersonic aircrafts, and nuclear technology). Phosphorus-containing copper-silver solder alloys are self-fluxing on copper, silver, and tin-bronze surfaces. Increasing silver content and decreasing phosphorus content improve the workability and electrical conductivity. These alloys find applications in electrical engineering and in cryotechnology, where solder connections have to withstand temperatures extending over wide ranges down to -50 °C. High purity silver and gold solder alloys are especially suited for hard solder processes under vacuum, protective gas, or in reducing atmospheres. The purity of the components is 99.99% minimum. They must also be free of metallic and nonmetallic constituents that have high vapor pressures at the soldering temperature (e.g., zinc, cadmium, arsenic). Palladium additions increase the corrosion resistance and the mechanical strength. These alloys are also used for joining metallized ceramic with metals for solder connections in vacuum tubes (e.g., magnetron, clystron, thyratron, and X-ray tubes), and for solder connections of molybdenum, tungsten, zirconium, and beryllium (Zimmermann, 1966). Silver-free, hard solder alloys based on palladium contain copper, manganese, and nickel. They have, compared to the silver solders, higher melting temperatures, improved strength, higher temperature stability up to ^800°C, and good corrosion resistance. They serve for brazing stainless steel, tool steel, and also refracto-

526

9 Noble Metals and Their Alloys

ry metals. Platinum-based solder alloys are used for solder connections which have to withstand very high temperatures. Table 9-42 gives some examples of commonly used noble metal based brazing alloys. For joining ceramic and metal, the ceramic surface first has to be metallized with an Mo-Mn preparation (coating with pastes, drying, and firing) and then galvanically plated with nickel. Direct soldering of ceramic with metal is possible with so-called "active solders", which contain ^ 2 - 4 % titanium or zirconium (Table 9-

Table 9-42. NM-containing brazing alloys. Type

Solidus Liquidus

Composition

Ag-Cu Ag40Cul9Zn21Cd20 Ag60Cu27Inl3 Ag44Cu30Zn26 Ag72Cu28

595 605 675 779 E a

630 710 735

Ag-Mn Ag85Mnl5

950

950

807 901 100

810 950 1120

Ag-Pd

Ag68.4Pd5Cu26.6 Ag54Pd25Cu21 Ag75Pd20Mn5

Pd-Cu Pd-Ni

Pdl8Cu72 Pd60Ni40

1080 1237

1090 1237

Au-Cu Au80Cu20 Au33Cu65 Au-Ni Au82Nil8

890 990 950

890 1010 950

E = Eutectic composition.

43). Bonding is performed by reaction of the refractory metal with the ceramic oxide (A12O3), forming an intermediate bond layer (Mitzuhara, 1987). 9.3.7.7 Dispersion-Hardened Alloys

The hardening effects due to solid-solution or precipitation hardening are usually limited. The alloys also exhibit decreasing strength with increasing temperature far below their melting points. A further disadvantage is caused by the marked decrease of the electrical conductivity, which is especially important for electrical contact material. Materials with much better mechanical properties at high temperatures and reasonable electrical conductivities can be produced by dispersion hardening. This is effected by incorporating in very fine dispersion, small amounts (0.1-2 vol.%) of very small particles of high melting compounds (oxides, carbides), which are insoluble or only become soluble near the melting temperature of the matrix. Dispersionstrengthened noble metal alloys are used for applications where high mechanical strength, corrosion resistance, and electrical conductivity are important. The strength of the dispersion-hardened metallic matrix depends of the size and the particle distance of the incorporated parti-

Table 9-43. NM-containing active solder alloys (Krappitz et al., 1991). Alloy

CB1 CB2 CB4 CB5 CB6 CS1 CS2

Composition (wt.%) Ag

Cu

72.5 96 70.5 64 98 86

19.5

Sn

In

Ti

5

3 4 3 1.5 1 4 4

26.5 34.5 1 10 4

Melting range

Brazing temperature

Pb

( C)

(°C)

92

730-760 970 780-805 770-810 950-960 221-300 320-325

850- 950 1000-1050 850- 950 850- 950 1030 850- 950 850- 950

527

9.3 Alloys

cles, which cannot be cut by dislocations. Strengthening is furthermore effected by precipitation of the additives at the grain boundaries. Dispersion-hardened noble metal alloys are produced by powdermetallurgical methods, supported by milling of the mixed powders (mechanical alloying), by internal oxidation of the base metal components in the noble metal alloy, or by coprecipitation of the noble metal powders and refractory oxides. The highest strengths are obtained by internal oxidation of alloy powders. The powders have to be densified and sintered to yield compact materials, and can then be worked by conventional manufacturing processes to rods, sheets, wires, or other product types. Dispersion-hardening by internal oxidation needs sufficient solubility of the oxygen in the base metal to obtain reasonable reaction times. This method is especially suitable for hardening silver alloys because of the high solubility of oxygen in silver. In the described concentration range, the particles have a spherical shape and are less than 20 nm in diameter. The particle size is, owing to the big difference in the energy of formation of the oxides of the matrix and the dispersoid (A12O3: 125 kcal/g atom), independent of the temperature and time of the oxidizing reaction (Poniatowski and Clasing, 1968). Figure 9-26 shows the hardness of dispersion-hardened silver compared to that of pure silver. Disperson-strengthened silver alloys are used for seal rings in oxygen-compressing equipment, for low voltage slide contacts, and corrosion resistant apparatus parts of chemical equipment. The solubility of oxygen in platinum is very low (< 1 ppm at 1400 °C). Internal oxidation of alloys based on platinum, which have been alloyed, for example, with 0.2 wt.% Cr or 0.8 wt.% Zr, is only possible after addi-

_160 E

^e ^120

-

hot formed

OJ

f 80 .E 40 CO

200

400 600 Temperature T

800

1000

Figure 9-26. Brinell hardness of (a) fine silver, (b) finegrain silver (99.7%), (c) silver (99%) dispersion hardened by oxide-metal powder, and (d) silver (99.5%) dispersion hardened by internal oxidation. (1 kp/ mm 2 ~450kPa.) (Degussa, 1967).

tional alloying with palladium or rhodium, which improve the oxygen solubility. Dispersion-strengthened platinum and platinum alloys are usually produced by coprecipitation of the metals and refractory oxides, for example, with ZrO 2 [ZGS = zirconium oxide grain-stabilized platinum] (Selman et al, 1974). Besides the application of oxides for dispersion hardening, platinum metals can also be effectively dispersion hardened by incorporating very small amounts of refractory carbides, for example, 0.04-0.08 wt.% TiC. The efficiency of the strengthening additive is related to a partial decomposition and reaction of the carbide with the matrix. The great advantage of this low concentration is that many properties like the ductility, the working characteristics, the electrical conductivity, and also the recrystallization behavior are not seriously impaired. Recrystallized structures with only a few grain boundaries normal to the applied tensile stress withstand long-term heating periods at temperatures at which material with a randomly oriented crystal structure

528

9 Noble Metals and Their Alloys

has already undergone considerable grain growth, without further recrystallization or grain growth. These structures can be produced with TiC dispersion-hardened platinum by recrystallization treatment at 1400°C after extensive cold deformation (area reduction of 95 %) (Darling, 1973, 6). Dispersion-hardened platinum is used in glass fiber manufacture, where platinumrhodium alloys cannot be used for melting optical glass because of discoloration effects. Dispersion-strengthened platinumrhodium alloys also serve as heater elements in high temperature furnaces. They combine high corrosion resistance and high mechanical strength at elevated temperatures. 9.3.7.8 Composite Materials Composite materials based on noble metals are primarily designed for electrical contact materials. Compared to dispersion-strengthened alloys, composite materials contain much higher proportions of either metallic or nonmetallic additives, which give rise to considerable changes in the material's properties. While dispersion-strengthened materials essentially show the characteristic physical and chemical properties of the matrix metal, the properties of composites are far more complex and are determined by the contributions of all the components: (1) combination properties comprising "sum properties" from additive effects of the single phases and "product properties" due to interactions of the involved components; and (2) structural properties determined by the shape and extension of the phases present (Stockel, 1979). Sum properties result from the addition of proportional properties of individual components like the electrical conductivity, density, specific heat, etc. Specific prod-

uct properties may arise from bimetal effects or from the combination of magnetostrictive phases with piezoelectric phases. In the field of electric contact materials, besides combination properties, the contribution of the structure is very important. Four basic structure types are distinguished: • • • •

particle composite materials, fibre composite materials, infiltrated composite materials, and laminated composite materials.

Particle composite materials are predominantly made from mixtures of silver powder with oxides, carbon, or nickel. AgCdO composites are made either by mixing metal and oxide powders, pressing, sintering, and extrusion, or by internal oxidation of alloy sheets or alloy powders followed

Ag

MeO Ag

(powder)\

/(powder) Mixing (milling)

Atomizing (powder)

Casting

Extrusion

Internal oxidation

Extrusion

Forming

PMP

Rolling

Contact

Extrusion

Internal oxidation

Forming

Forming

I

I

Contact Contact Figure 9-27. Production processes for Ag CdO composites; Me = base metal, PMP = powder-metallurgical processing.

9.3 Alloys

'&&#*•;• •±f&.*r~ ~siiiz£&-

529

(b)

') '\*J;; • ,<:'- •*••**:! • >"•' :•. >: :..-"J' .*: •;' V.;1!- ' l k : ^ i ^ S ^ ^ ^ ^ ^ ^ » - ^ . ^ i f f i s f c ' *. (d)

*^S1" ' r s - ^ v r ^ i ^ f s i^^'-SSr^^'t^^:-^'''-"

mm I

'

' "^"•-'*''-***^^T;t^S^^^S» '' ***• is.'' * •• -*^" —

by powder-metallurgical densification and sintering (Fig. 9-27). Fiber composite materials with oriented fibers are fabricated by the extrusion of powder mixtures (Fig. 9-28). They are preferably used as contact material in silver-metal oxide and silver-graphite type electrical contacts. Contact faces that have been cut vertical to the axis of the fibers have increased erosion resistance in makebreak contacts compared with composites with random orientation of the composite particles. Silver composites with oriented fibers are alternatively produced by placing thin nickel wires inside silver tubes, which are then drawn together, cut, made to bundles, and usually severely deformed by rolling, forging, or drawing. By repeating these operations, composites with fiber diameters down to 50 nm can be produced.

-•-..«,-;.,s.

Figure 9-28. Fiber structure of extruded powder composites: (a) cross-section and (b) longitudinal section of extruded Ag-C composite, (c) cross-section and (d) longitudinal section of extruded Ag-Ni composite (Degussa).

In a similar way tube-shaped inclusions can be formed inside the silver matrix. This method allows the production of silver-nickel composites, with high nickel contents, that have improved technological properties compared to tungsten in various kinds of switch gears (Fig. 9-29). Silver-nickel fiber composites are magnetic. High coercivity results if the nickel fiber diameters become so small that no Bloch wall can be generated, and therefore only one single domain exists. Figure 9-30 shows the dependence of the coercive force on the fiber diameter. Coercive values of 526 A/cm at 77 K and 333 A/cm at 295 K have been obtained (Stockel, 1974). Infiltrated composite materials are produced by first forming a skeleton of high melting metals like tungsten, molybdenum, or tungsten carbide, into which liquid

530

9 Noble Metals and Their Alloys

silver is infiltrated, penetrating by capillary forces. They show superior behavior in arc erosion resistance because the evaporating silver prevents a more pronounced increase of the arc temperature and thereby protects the refractory metal from melting. Silver-infiltrated materials with 60-80 wt.% tungsten withstand arc erosion better than pure tungsten. Rocket nozzles made of silver-tungsten composites can withstand softening by the heat that develops during reentry into the atmosphere by friction, because the cooling effect of the melting and evaporating silver prevents softening or melting of the tungsten skeleton. Laminated composite materials are made by common methods like hot pressing, diffusion bonding, hot rolling, cold rolling, or explosive welding.

9.4 Manufacturing: Materials and Methods 9.4.1 General The manufacture of noble metals and alloys is predominantly done by common metallurgical methods: melting, casting, sintering, mechanical working, separating, and joining. The surface layers are either clad or deposited by galvanic, sputtering, evaporation, or chemical vapor deposition (CVD) processes. The required mechanical properties are adjusted by controlled processing and by thermal treatment. Basically three main types of production materials have to be considered: 1. Bulk material made by melting. 2. Powders and powder compacts. 3. Components and compounds for the production of coatings. Figure 9-31 gives a general outline of the materials, methods, and their functions. 9.4.2 Bulk Material Made by Melting 9.4.2.1 Melting Conditions Gold, silver, and their alloys are melted in graphite, clay-graphite, and SiC crucibles, platinum and PGMs, except iridium, in A12O3, MgO, or ZrO 2

Figure 9-29. Ag-Ni wire fiber composite rivet: a) manufacturing, b) cross section (Stockel, 1974).

531

9.4 Manufacturing: Materials and Methods

copper crucibles, platinum and PGMs under a protective gas atmosphere. High temperature melting offers the chance to purify the material by the evaporation of accompanying impurities of higher vapor pressure ("consolidation melt"). Table 944 shows the purification effect of melting iridium at 2700 °C. 1.2 1.6 r(um) Figure 9-30. Coercivity of Ag/Ni wire fiber composites with small fiber diameters. Ag/Ni20; x 77 K, A 295 K (Stockel, 1979, 2).

Table 9-44. Purification of iridium by consolidation melting at 2700 °C (Baving et al., 1991). Impurities

crucibles, and usually heated in induction furnaces. For melting at temperatures above 2000 °C, the methods of vacuum-arc and electron-beam melting are applied. The material is melted in water-cooled

Melt

Ingot

Rods/wires

Sheet/band foil

Punched parts

Clad materials

From pieces

Anodes

Components Chemical equipment Electrical engineering • Electronic Catalysis Jewelry Gauzes Contacts PVD targets Evap. slugs Solders Bond wires

Sputter targets

Jewelry Dentistry

Sintered compacts

Composites Electronic contacts

Pastes

Lacquer

Adhesives

Before

After

120 400 700 50 25

< 5 <10 <50 <10 < 5

Compounds

Powder

*—

Profiles

Ag Cu Te Cr Mg

Concentration of impurities

>

Conductors Resistors Electrodes Electrical devices Circuits Assembling Solders

Suspension, Single piece production injection

Salts

Photo CVD

Surface coating Conductor Corrosion Protection Decor Solution > Jewelry for thermal Catalysis decomposition Electrical engineering Volatile Electronic compounds CVD Galvanic baths

Solution for chemical reduction

Heterogeneous catalysis

Complex organic compounds

Homogeneous catalysis

Figure 9-31. Overview of NM-containing production materials.

532

9 Noble Metals and Their Alloys

9.4.2.2 Casting Casting is either done discontinuously, producing ingots, or continuously, producing endless, mostly simple, profileshaped rods, bands, or strips. Discontinuous casting is carried out in water-cooled copper molds. If concentration inhomogeneities are formed during solidification, the cast ingots have to be heat-treated for homogenization. In the process of continuous casting, solidification takes place when the liquid metal passes an extended cooling section of defined length and controlled decreasing temperature (Fig. 9-32). This method is cost-advantageous because it minimizes the necessary subsequent mechanical work. It produces consistent material quality along the whole profile length. The faster solidification which continuously occurs in limited, small volume parts prevents the formation of undesirable segregation. On melting in air, the process of continuous casting produces material with a low oxygen content ( ^ 10 ppm for pure silver and ^30-50 ppm for silver alloys). Horizontal continuous casting is particularly used in the production of large quantities and material of large cross section. It is the dominant casting process of silver and silver alloy production technology. Industrial plants reach production capacities of 2000 t/year. Profiles of dimen-

sions ^330x500 mm are produced with speeds of up to 100 kg/h. Smaller equipment works using vertical continuous casting. This is preferably used for high quality material, e.g., gold alloys in kilogram quantities. The dimensions of the strips are of the order of 5 x 80 mm, rod or wire diameters from 10 to 50 mm. Investment Casting Small items of more complicated shape are cast by the investment casting method, preferably using molds which have been made by the "lost wax" process. In this process, wax models of the desired shape are formed in a split mold, then embedded with a hardenable paste, based on quartz. After hardening the wax is melted out, the mold fired, and the noble metal poured in. The cooled metal piece is removed and brought to a final finish by polishing. 9.4.2.3 Mechanical Processing The first operation after discontinuous casting is in most cases extrusion in presses with intermediate heating. Subsequent forming processes include forging, hammering, rolling, drawing, cladding, cutting, bending, punching, and deep drawing. Pure palladium, silver, platinum, and gold can be cold rolled down to 99 % area reduction without intermediate heating.

Figure 9-32. Continuous casting of silver (schematic diagram): (a) pouring ladle with molten metal, (b) holding furnace, (c) continuous water-cooled mold, (d) rollers, (e) circular saw, (f) milling machine, (g) coiling machine (Ullmanrts, 1993).

9.4 Manufacturing: Materials and Methods

Cold-deformed noble metal alloys have to be annealed to soften them for further processing. Thin foils are rolled in a package arrangement. The very thin gold leaf used for mechanical gold-plating is well-known. Production starts with sheets of «20 (im thickness. These are placed between sheets of vellum, beaten down, and finally hammered by hand, using separating membranes of the cecum of cattle ("goldbeater skins") to a thickness of ^0.2 mm (Nicholson, 1979). A modern method of making gold leaf is cathode sputtering. Extruded rods are transformed to wires via classical wire-drawing technology: drawing through dies with decreasing diameters in defined reduction steps. Pure gold and pure silver wires can be drawn to very small diameters without intermediate annealing. Final annealing is done to adjust the mechanical properties (e.g., tensile strength and elongation) according to application requirements. Wires of platinum, palladium, and gold with very small diameters of ^ 1 - 1 0 |im ("Wollaston wires") are drawn in an outer envelope of silver. Gold wire made from high-purity gold can be drawn down to w 10 Jim without an outer cover. Pure gold wire of 25 jim diameter with specified values of tensile strength and elongation is used as "bonding wire" for connections in electronic devices (see Sect. 9.5.2.2). Cutting processes include turning, drilling, and punching. To minimize material losses, profiled strips and precision tools and machines are used. The main technical items are spin nozzles, contact pieces, and punched preform materials for assembling, soldering, or brazing devices and equipment. Noble metal specific procedures are engraving and guilloching of silverware, punching and glaze cutting of jewelry items made from platinum and white gold, punching of holes in spin noz-

533

zles, and welding of noble metal layers on base-metal supports for electrical contact materials. Joining is performed by soldering and welding with suitable alloys. Cladding is done by special cold- or hotrolling operations. It is of special interest for the production of mechanically resistant coatings on base-metal substrates. For better adherence, cladding is often combined with the application of an intermediate metallic layer, which may be pressed, welded, or rolled together with the outer layer on the base-metal surface. These processes play an important role both for jewelry and technical applications (for example, "gold-double" fabrication and material for electrical contacts). Copper alloy foils, which are partially clad with silver and gold strips in "inlay" or "toplay" design are precursor materials for punched electrical contact rivets and semiconductor lead-frame components. 9.4.2.4 Thermal Treatment

Thermal treatment during the manufacture of noble metals on semifinished and finished workpieces serves, besides giving stress-free annealing, different purposes: a) Lowering of strength by recovery or recrystallization: Annealing temperatures and times depend on the alloy composition, the degree of cold working, and the purity of the material. Pure silver and oxide-containing silver composites like AgCdO, AgSnO2, or some silver alloys with high contents of gold, palladium, or platinum can be heated in air. All other noble metals and alloys must be annealed in protective gas atmospheres. Small pieces can be treated in salt melts of alkali chlorides, carbonates, or cyanides. Oxygen or oxide-containing materials, e.g., Pt/ThO 2 , Ag/CdO, Ag-Cu alloys containing CuO,

534

9 Noble Metals and Their Alloys

and silver pieces which have been previously annealed in air, may not be annealed in a hydrogen atmosphere. In these cases, the hydrogen reacts with the dissolved oxygen, forming either local inhomogeneities by the formation of alloys or blisters of water vapor ("hydrogen pest"). At higher temperatures and after long-term heating in flowing gas, evaporation occurs. Ruthenium, silver, osmium, iridium, and platinum show more pronounced vaporization losses if they are annealed in the presence of oxygen (instead of argon or nitrogen), because the volatile oxides that form are removed by the passing gas stream. b) Strengthening by internal oxidation of solute base-metal components (dispersion hardening): On internal oxidation, dissolved oxygen reacts with the alloyed base-metal atoms. Fine, dispersed oxides are formed with particle sizes of ^ 1 jLim. Materials show higher hardness and lower ductility at room temperature, but have a considerably higher mechanical strength at high temperatures, and higher electrical conductivity. c) Homogenization of castings (solution annealing and high-temperature heating): Cast ingots are homogenized by annealing to equalize concentration inhomogeneities. Heavy segregations can only be homogenized by solution heating at high temperatures. d) To obtain hardening by precipitations or ordered structures: The following annealing temperatures are used: • • • •

Silver alloys 200-400 °C. Cold alloys 400-600 °C. Platinum alloys 500-1000 °C. Palladium alloys 600-1100 °C.

Annealing is usually done in air or in salt baths. e) Sintering of powder compacts and infiltration of sintered refractory-metals'

skeletons with liquid silver: Sintering of pressed or loose-poured noble metal powders begins at half the melting temperature (measured in K). At this temperature, sintered bodies with densities between 25 and 50% of the density of the bulk material are formed. Higher densities are reached with higher temperatures and subsequent deformation (rolling, extruding) at elevated temperatures. A slow increase in the sintering temperature is needed to enable adsorbed gases to escape. Sintering processes may only be carried out in hydrogen atmospheres if the powders are free of oxygen. 9.4.3 Powders and Powder-Metallurgical Processing 9.4.3.1 Powders Noble metal powders are needed to produce sintered bodies, as constituents in metallizing and soldering pastes, and as catalysts. The powders are characterized by a number of different parameters, e.g., particle shape, particle size, particle size distribution; specific surface [measured by adsorption of nitrogen by the BET (BrunauerEmmett-Teller) method (Brunauer et al. 1938)], and tap density [loose poured powder is densified by vibration, standard method determined by ASTM 527-85] (Stanley-Wood, 1977). These parameters influence the working behavior during pressing and sintering (density, sintering temperatures), the chemical surface activities and the rheology of noble metal powder containing pastes. Powder Production The main powder production technologies are: chemical reduction, electrolytic deposition, and atomization of metal melts, leading to characteristic, different properties.

9.4 Manufacturing: Materials and Methods

Chemical reduction works with noble metal salts, predominantly in aqueous solution. Additions of protective colloids and wetting promoting compounds enable control of the particle size distribution and prevent agglomeration during powder deposition. Reducing agents include solutions of sodium formate, hypophosphoric acid, or hydrazine. The powders are very fine and have irregular or spherical shapes. Particle sizes depend on the process conditions. They are usually in the range of 0.520 jam. These powders are preferably used for metallizing and electrode pastes. Extremely fine powders ("metalblacks") of rhodium, palladium, iridium, and platinum with BET surfaces between 9 and 38 m2/g are produced by the chemical reduction of alkaline solutions of the hexachloro-complex compounds with formaldehyde. Electrolytic deposition of powder is predominantly used for powders used in the production of sintered contact materials. Deposition proceeds continuously in an electrolytic cell. The noble metal anode dissolves under the action of the current, and the powder crystals are deposited at the cathode, from where they are mechanically removed. The powders are of dendritic shape with particle sizes of up to 200 jim in diameter. Atomizing from the melt is primarily used for the production of powders for solder pastes, for sinter processes, and for high purity materials for filling preparations in dentistry. Atomizing is done by mechanical dispersion (for instance with rotating disks) or by spraying through nozzles and quenching the particles in water or a cold gas stream. High purity spherical powders with good flow properties are formed in inert gas atmospheres.

535

Types of Powder Lamellar-shaped powders (flakes) are obtained by milling the deposited powders in ball or attritor mills in liquid suspensions. Plate thicknesses are in the range of 0.1-1 jiim. The flakes are used as pigments in conductive lacquers and conducting adhesives, the face contact between the flat grains leading to increased conductivity. Nanocrystalline noble metal powders with particle sizes in the range of 101000 nm and colloid dispersions (especially from gold and silver) are produced by the condensation of evaporated materials (for instance by electron-beam evaporation) on cool surfaces, by generating an arc between noble metal electrodes immersed in suitable liquids, or by chemical deposition out of a solution that contains protective colloids. As a result of the increased ratio of surface atoms to bulk atoms in each grain, nanocrystalline powders have a stronger catalytic action and lower sintering temperatures. Compact, densified materials made from nanocrystalline powders have higher hardness and higher creep resistance compared to conventionally produced noble metal bulk materials. Ultrafine silver powder in an organic binder material forms films with very stable electrical resistance; these are used as sensing tape for the detection of the position of magnetic tapes (Benner et al., 1991, 2). Uses Silver powder finds the widest application. Important quantities are used in electronic pastes for the production of conductive layers and electrodes, often in combination with palladium or platinum. Powders of silver-containing hard solder and soft solder alloys are constituents of solder pastes.

536

9 Noble Metals and Their Alloys

Powder-metallurgical processes are used for the production of contact materials (Ag/CdO and Ag/SnO2) and electrodes for batteries. Elemental silver powder is used for catalytic actions. Fine silver powder is produced in an aqueous solution of AgNO 3 with NaCOOH as the reducing agent 2AgNO 3 + NaCOOH -* 2 Ag + CO 2 + NaNO 3 + HNO 3 (9-17) High purity silver powder is required for metallizing and electrode pastes in electronics. Production is done by a two-step process + Na 2 CO 3 -+Ag 2 CO 3 + 2NaNO 3 (9-18) II) Ag 2 CO 3 + HCHO -> 2 Ag + 2 CO 2 + H2 (9-19) which allows cleaning alkali away before reduction of the silver carbonate with formaldehyde. Coarse-grained silver powder grades are formed by precipitation out of aqueous silver nitrate on copper surfaces Cu + 2 AgNO 3 -> Cu(NO 3 ) 2 + 2 Ag (9-20) or by electrolytic dissolution of silver anodes in aqueous silver-salt solutions or

aqueous silver oxide suspensions, and electrolytic deposition of silver crystals. Palladium powder is produced by chemical reduction of aqueous Pd(NO 3 ) 2 or PdCl2 solution. The majority of this powder finds pure, alloyed, or mixed with silver powder (mainly of composition Ag-Pd 30 wt.%) application as a constituents of metallizing pastes for the production of inner electrodes of multilayer capacitors. Special, fine particle sizes are needed to get sufficiently thin and homogeneous layer structures. Alloyed powders are produced either as composite blends or by coprecipitation from mixed-salt solutions, followed by thermal annealing of the dried powder. Table 9-45 gives some characteristic properties of silver, palladium, and Ag-Pd powder grades. Gold powder is produced by the chemical reduction of tetrachloro-gold acid (HAuCl4). Depending on the reducing agents, either spherical particles (between 1 |im and 10 jim) or plate-like-shaped particles (between 0.5 jim and 50 |nm) are formed (Fig. 9-33). Gold powder is used in preparations for die bonding, for the production of bond-pads, and in thick film metallizing pastes to produce conductor layers. Gold conductor layers are used in-

Table 9-45. Typical silver, palladium, and Ag-Pd powder grades. Metal

Manufacturing method

Grain shape

Grain size (Mm)

Tap density (g/cm3)

Specific surfacea (m2/g)

Ag

chemical reduction electrolytic deposition

microcrystalline dendritic

0.5-2.0 1-200

0.8-5.0 4.5-4

0.1-2.0 0.1-0.5

2 - 40

2.5-5

0.2-1.8

<20 <1.2

0.8 4.0

2.5 2.3

<1.2

3.3

2.2

<9

3.4

3.3

1

ball milling Pd Ag-Pd30

chemical reduction coprecipitation

1 flakes crystalline spheres spheres 1

i flakes BET in N 2 adsorption.

9.4 Manufacturing: Materials and Methods

537

(d) Figure 9-33. (a) Powder of gold-based dental alloy made by gas-spraying, mean grain size 20-40 urn (Baving et al, 1991, 2). (b) Powder of silver-copper solder alloy made by water-spraying, mean grain size 30-60 jam (Baving et al., 1991, 3). (c) Palladium powder made by chemical reduction of aqueous solution, grain size »0.2 urn (Dorbath et al., 1991, 3). (d) Gold powder with platelet shaped particles (0.5-50 urn), made by chemical reduction of aqueous solution with special additives (Dorbath et al., 1991, 4).

stead of conductor layers made from silver in highly reliable electronic circuits, where silver migration, which can occur in certain applications under current, may cause problems. Preparations for the decoration of porcelain and ceramics, and for the filling of cavities in dentistry, are further application fields for different gold-powder grades. Platinum powder serves in silver, silverpalladium, or gold-containing thick film pastes to increase the resistance against scavenging during soldering operations with tin and Sn-Pb solder alloys. 9.4.3.2 Powder-Metallurgical Processing

I |

Powder-metallurgical processes are used for compacting hard and difficult-to-work noble metals or alloys (rhodium, iridium, platinum-cobalt and platinum-cobaltchromium alloys), to produce composite contact materials (Ag/Ni, Ag/CdO, Ag/ SnO2, Ag/C, Ag/Al 2 O 3 , W-Ag, Mo-Ag)

538

9 Noble Metals and Their Alloys

and dispersion-hardened materials for construction, for production equipment, and applications in electrotechnical engineering (Pt/ZrO 2 , PtRh/ThO 2 , Pt/TiC). High-gold-containing jewelry mass articles like chain links, bracelets, etc., are made from powder suspensions by injection molding and subsequent sintering. The densities of pressed and sintered materials are influenced by particle size, shape, applied pressure, and temperature. The sintering of pressed or loose-poured noble metal powders begins at half the melting temperature (measured in K). Applied pressures for silver powders are around 0.1 MPa, for platinum metals 0.20.3 MPa. Sintered bodies reach densities above 90% of the theoretical density. Further densification can be done by subsequent hot pressing, hot rolling, or extrusion up to ^ 9 8 % . Ready-formed parts such as electrical contact pieces are pressed in shaped dies by pistons. Using different powder compositions two layers can also be pressed together (for example, AgCdO with an outer layer of pure silver). Special contact materials consist of sintered porous skeletons of tungsten or molybdenum which are subsequently impregnated with liquid silver. The production of larger-size ingots is performed by cold isostatic pressing (CIP) or hot isostatic pressing (HIP). Powders or powder mixtures are filled into plastic forms and introduced into a closed, highpressure steel container. In CIP, homogeneous pressure transfer is accomplished using liquids; in HIP the pressure transfer takes place via filling the container with gas. An applied pressure of 6000 bar (600 MN m~ 2 ) is used for CIP, and 2000 bar (200 MN m~ 2 ) for HIP with a temperature of up to 2000 °C. The compacted ingots have the dimensions of 150 mm diameter and up to 500 mm length. Hot iso-

static pressing is also used for additional densification of porous, cast or injectionmolded form pieces. A special, continuous powder-pressing process is used for the direct compacting of powder to wires of silver solder compositions (Green, 1972). 9.4.4 Textures 9.4.4.1 General Classification In polycrystals the specific grains have different orientations with respect to the geometry of the bulk material. The distribution of the orientations is called the texture. In a more special sense, the preferred orientations deviating from the statistical distribution are defined as the texture. In a wider sense, the oriented growth of single crystals out of vapor, melts, or galvanic solutions can also be treated under this topic. Textures are formed in polycrystalline materials during solidification from melts, by mechanical deformation (deformation textures caused by rolling, extrusion, drawing), and by thermal treatment of solid polycrystalline materials (recrystallization textures). If the basic deformation mechanisms of different materials are the same, differences can exist between the deformation textures, for example, for silver and gold which form different proportions of <111> and <100> fiber textures in drawn wires. In addition to the basic properties, e.g., crystal structure, stacking fault energy, etc., the texture formed is affected by the manufacturing process conditions, by the starting orientation, and, in the case of the deposition of surface layers, by the orientation of the crystals of the substrate surface (epitaxy). As most single crystals behave anisotropically in their physical and mechanical properties (for example, the elasticity of gold crystals in different crystal directions), textures can have a marked influ-

9.4 Manufacturing: Materials and Methods

ence on polycrystalline material properties. The manufacture of special textures in polycrystalline materials by controlled production steps can serve to suppress or to generate intended anisotropic behavior ("texture tailoring"). Texture formation in noble metals and alloys has been examined especially for silver, gold, and to a lesser extent, palladium and platinum. 9.4.4.2 Single Crystals Silver whiskers, formed on the contact of silver with silver sulfides, selenides, or halogenides have a <112> orientation along the fiber axis. Galvanic deposits of silver show different orientations depending on the operating conditions. Low current density affects the formation of <111 > and <001> textures parallel to the current line direction. With higher current densities, a random distribution is formed. On {111} oriented gold surfaces, silver is also deposited in a {111} texture. Similar behavior is found for galvanically deposited gold layers. Thin, galvanic gold coatings show the orientation of the substrate surface, whereas thicker layers form, depending on the concentration of the solution and the current density, their own structures, preferably with {111} and {211} parallel to the substrate surface. In the vapor deposition of gold, various different textures can also be formed, depending on special crystal orientations of the substrate surface, the vapor concentration, and the deposition temperature (Raub, 1940, 2). Gold and silver crystals, formed by the solidification of melts either in the form of dendrites or columnar structures, have the <100> direction along the dendrite axes (Wassermann and Grewen, 1962).

539

Deformation Textures Single crystals of silver and gold, if drawn without a matrix, form, independent of the starting orientation, a <211> orientation parallel to the drawing direction. If the starting orientation is <111> and the material is drawn through dies, twinning occurs at low deformation, forming a <100> fiber texture. Further drawing to higher deformation leads to a reduction of the <100> orientation in favour of <111> again. This transformation begins for gold at a lower deformation, so that under the same conditions in gold crystals, compared to silver, a much greater proportion of <111 > is formed (Ahlborn, 1965).

9.4.4.3 Polycrystalline Materials Deformation Textures In cold-deformation processes of f.c.c. metals, the slip planes are the four {111} planes, with three <110> slip directions in each slip plane. The {111} planes are also the twinning planes. Extruded silver rods have a predominantly <100> texture with a proportion of <111> and <543>. The textures of cold-rolled silver and its alloys with gold and copper are the same as that of brass with {112} orientation as main orientation. Gold has additional {111} orientation as a second orientation in the same direction. The textures of rolled silver and silver alloy sheets cause anisotropy of the tensile strength (Schmid and Wassermann, 1925), which has its lowest value in the rolling direction and a maximum value in the cross direction; this influences the deformation behavior in deepdrawing operations. This effect can be suppressed by repeated rolling in alternating directions (cross rolling). Very thin, beaten gold foils have orientations with the {100}

540

9 Noble Metals and Their Alloys

and the {110} planes in the foil plane (Regenet and Stiiwe, 1963; Raub, 1940). Wire drawing of silver, gold, and palladium results in a double-fiber texture in the <111> and <100> directions. For platinum, only the <111> fiber texture was detected. Polycrystalline gold and silver wire behave likewise, as described for the single crystals. At high deformation, gold wire forms much larger amounts of <111> texture, where for silver deformation twinning is responsible for the predominant formation of <100>. Twinning takes place in gold wires if they are drawn at -190°C, forming up to ^ 5 0 % <100) orientation (Ahlborn and Wassermann, 1963). Recrystallization Textures Recrystallization textures arise from cold-worked, textured materials by annealing at temperatures above the temperature of recrystallization. Altered textures originate out of the existing structures if new crystals form with preferred growth along special planes, and are therefore related to the textures which were previously formed by cold deformation. By annealing cold-rolled silver and silver-copper alloys at up to 750 °C, a new crystal orientation is formed with the {113} plane parallel to the rolling plane. Consequently, the tensile strength of the recrystallized material has a higher value in the rolling direction and a lower value in the cross direction. Annealed sheets that have been cross rolled in alternating directions show a minimum tensile strength at 45° to the rolling direction. The <100> fiber texture of hard-drawn silver wires is transformed by annealing to a <211> fiber texture by forming extended crystals in the direction of the wire axes. The annealing of cold-drawn gold wires that have a predominantly <111 > fiber tex-

ture, at temperatures of 200-300°C, changes the texture from <111> to <100>. This leads to a marked decrease of the modulus of elasticity both in the direction and normal to the direction of the fiber axis, and effects typical changes in the stress-strain behavior (Matucha, 1982). The Debye-Scherrer photograph of a pure gold wire of 25 jim diameter after a short annealing time shows the double fiber texture with <111> and <100> in the direction of the fiber axis. (Fig. 9-34). 9.4.5 Deposition of Coatings Noble metals and noble metal alloys are used for numerous different functions as surface layers on metal, ceramic, and plastic material supports. Besides jewelry and decoration, the main technical applications of noble metal coatings include electrical conductors, resistors, contacts, fuse elements, mechanical connections, and corrosion resistant, optical, and heat-reflecting components. The surface layers are produced: a) from bulk material components by mechanical cladding (rolling, pressing), thermal (soldering, welding) cladding, and physical vapor deposition (PVD: evaporation, and sputtering), or b) from chemical compounds or solutions by galvanic deposition, chemical reduction, noble metal preparations (lacquers and pastes), thick film pastes (screen printing and firing), thermal decomposition (e.g., resinates), and chemical vapor deposition (CVD). 9.4.5.1 Coatings from Bulk Material Components Mechanical and thermal supported cladding processes are preferably applied for the deposition of thicker coatings, as

9.4 Manufacturing: Materials a n d Methods

(222) (400)

(113)

(202)

(111) (200)

(111)

541

(202)

(200)

a ooo>

113

Figure 9-34. (a) Debye-Scherrer photograph (Straumanis method) of gold bond wire of 25 um diameter, showing double fiber texture in <100> and <11 Indirections along the wire axes (Matucha, 1982). (b) The theoretical diagram of diffraction lines shows, for comparison, the position of the enforced reflexes corresponding to the <100> • and <111 > A fiber orientations.

used for electrical contacts and for jewelry items ("double" coating). Gold coatings made by rolling are fabricated by first soldering gold sheets, usually in 14 carat grades, on blocks of stainless steel, copper, or bronze. These are then rolled or drawn to the desired shapes, producing strips and wires with a dense and adherent gold coating, usually of 5-10 jam thickness. The clad material is mainly used for spectacle frames, jewelry pieces, and watches. "Fire gilding" has been used to plate handmade articles and for the restoration of architectural objects, using gold amalgam applied on the surfaces and

heated to evaporate the mercury. Because of the strong toxicity of the mercury vapor, this process is no longer in use. Gold leaf, 0.1-1 jam thick, is used to gild wooden frames and book edges. PVD coating by sputtering or evaporation is mainly used for thin, high purity coatings for electrical contact pads in electronic devices and thin film circuits, for thin, transparent, energy-saving noble metal (preferably silver) layers on architectural glasses and for optical mirrors (silver plus rhodium). Platinum and iridium can also be deposited on substrates by vacuum evaporation, sputtering, or ion plating.

542

9 Noble Metals and Their Alloys

Evaporation slugs and sputter targets are materials for the production of PVD coatings (Schlott, 1995). 9.4.5.2 Coatings from Chemical Compounds or Solutions

Galvanotechnical processes find their widest application in jewelry and technical applications (Both et al., 1991). Coating is predominantly done in aqueous solution. Anodes for the deposition of silver coatings consist of silver which dissolves during the course of electrodeposition. Gold and platinum metals are deposited from the electrolyte bath using anodes made from titanium which has been coated with a platinum surface layer. Galvanic baths contain additives which serve for wetting, as a brightener, and as grain-refining agents. Additions of iron, nickel, or cobalt increase the hardness of galvanically deposited gold coatings. Silver and gold are commonly plated from cyanide solutions, the platinum metals from solutions of complex inorganic compounds. The coating thickness ranges from 0.1 to 20 \xm. Galvanic gold coatings have found widespread applications in the decorative (watch, spectacles, cutlery, writing implements, jewelry) field. The electrotechnical and electronic industries utilize the resistance to corrosion and abrasion, and the electrical conductivity for contact material and conductor components. The excellent soldering and bonding properties of pure, fine gold layers are used in the manufacture of semiconductor elements. The process of electroforming is used to manufacture hollow jewelry. Controlled electrolytic process conditions and special gold electrolytes allow the deposition of crack-free gold alloy coatings (usually produced in 8, 14, or 18 carat) on carrier substrates in thicknesses of up to 200 |im. The

hollow article is achieved at the end of the process by removing the substrate, which serves only as a mandrel. It consists either of a metallized wax or base metal, which can subsequently be removed by melting, dissolving, or etching. This method is especially suitable for producing very fine and exact replicas of surface structures (Mohan, 1976). Electrolytic deposition of silver is applied in industrial processes and in the manufacture of jewelry. The electrolytes used are solutions of cyanidic potassiumsilver salts; soluble silver anodes replace the deposited metal. Cu-Ni-Zn substrates ("German silver") and copper or nickel substrates can be electroplated directly, while iron needs an initial coat of either copper or nickel before silver-plating. Silver electroplating of plastic or ceramic materials also requires an initial undercoat of copper or nickel. The quality of silver-plating on cutlery is quoted in grams of silver per 24 dm 2 surface. For 90 g silver-plating (hallmarked "90"), the minimum coating thickness is 36 jim. Silver coatings on jewelry items have thicknesses from 5 to 10 |im, and plating thicknesses for electronic applications are in the range of 0.1-3 jum. Electroless silver plating is done by chemical reduction of ammoniacal silver nitrate solution with formaldehyde and sodium hydroxide as reducing agents. This method is applied for the coating of mirrors, heat-insulating glass vacuum containers, and for decoration purposes. Platinum, palladium, and ruthenium can be electrolytically deposited from aqueous solutions of complex inorganic compounds., Rhodium can be deposited from aqueous rhodium sulfate or rhodium phosphate solutions. Among the galvanically deposited PGM coatings, rhodium coatings have the highest hardness and op-

9.4 Manufacturing: Materials and Methods

tical reflectivity. Platinum and iridium can also be deposited by melt electrolysis of cyanide compounds at temperatures of 500-600 °C. This process is used for the electroplating of platinum and iridium on titanium, niobium, tantalum, and molybdenum. Applications include, for example, insoluble anodes for electrochemical processes, components for cathodic corrosion protection, and platinum-coated molybdenum leads of halogen incandescent bulbs (Simon, 1992; Simon and Zilske, 1995). Wear-resistant, galvanically deposited noble metal coatings can be produced by dispersing fine particles of A12O3, SiC, WC, or other refractory materials in the plating baths. Noble metal preparations contain the noble metal powders in suspension or mixture with liquid organic vehicles, i.e., special, soluble organic compounds (e.g., ethylcellulose) to adjust the viscosity and flow properties and inorganic components [e.g., finely dispersed glass or quartz powder ("frits"), base metal oxides] as bonding agents to ceramic substrates. These preparations are applied by brushing or screen printing onto the surfaces and then fired, forming dense, sintered, adherent noble metal coatings. They serve for metallizing ceramic components, electronic vacuum tubes, passive electronic devices, or thick film circuitry. Metallizing lacquers contain the noble metal powders in a suspension of organic lacquers without inorganic components. They can be hardened at room temperature or slightly elevated temperatures. The main applications for these lacquers are in the deposition of electrically conductive silver or gold coatings on plastics, sensor foils, and as the base metallizing layer for subsequent galvanic deposition of noble metal coatings. Liquid nobel metal preparations are based on nobel metal resinates. They serve

543

for the precipitation of very thin (<0.1 jim) noble metal layers of gold, silver, platinum, and palladium on ceramic, porcelain, glass, quartz, and mica, both for decoration ("liquid bright gold") and technological purposes. Preparations contain organic noble metal compounds (for gold: sulforesinates, for silver and platinum: organic thio-compounds or carboxylates). They are applied by painting, brushing, screen, or offset printing, and firing. Bright noble metal coatings are formed in air at temperatures between 500 and 1250 °C. Their adherence to ceramics and porcelain is effected by additions of special oxides (e.g., Bi 2 O 3 , SnO2) or oxide-forming compounds. Rhodium oxide prevents grain growth and agglomeration of the deposited noble metal pigments (Hunt, 1979). Platinum resinates are employed for decoration and for the production of conductive coatings on ceramic substrates onto which further layers can be deposited electrolytically. Thicker deposits (1-10 jim) are obtained by the addition of fine noble metal powders, which leads to the formation of matt surface ("burnished gold", "burnished silver") (Frembs, 1995). Ruthenium oxide (RuO2) coats are applied on titanium from anodes in chloroalkali electrolysis. Coating is performed by spreading an alcoholic solution of H 3 RuCl 6 onto the substrate, followed by heating in an air stream. Chemical vapor deposition (CVD) offers the possibility of producing coatings of all noble metals at moderate temperatures and of high purity at reasonable reaction rates. Precursor materials consist of complex metal-organic and inorganic noble metal compounds which can be vaporized and led by carrier gas stream to the substrate surface. Coatings are formed at elevated temperatures either by chemical reduction or by thermal decomposition on

544

9 Noble Metals and Their Alloys

Table 9-46. Main application areas for NMs. Metal

Application area

Compound

Application area

Ag

photography investment, jewelry, coins silverware electrical technology electronics brazing/soldering catalysis dentistry heat mirrors chemical technology

AgNO 3 Ag2O K[Ag(CN)2]

photography batteries galvano-technology

Au

jewelry, coins, investment electronics dentistry brazing/soldering corrosion protection reaction equipment apparatus components

K[Au(CN)2]

galvano-technology

Pt

catalysis jewelry, investment electrochemistry electronics sensors reaction equipment apparatus components glass industry autocatalysis

H 2 PtCl 6 PtNO 3

metal deposition on supports for catalysis

Pd

electronics electrotechnical industry catalysis dentistry jewelry autocatalysis

PdCl2 H 2 PdCl 4 Pd(NO 3 ) 2

metal deposition on supports for catalysis

Rh

catalysis electrochemical industry eletrotechnical technology sensors autocatalysis

RhCl3 Rh(NO 3 ) 3

metal deposition on supports for catalysis precursor material for CVD, etc.

Ru

electronics electrochemical industry catalysis

RuCl3

Ir

electrochemical industry petroleum chemistry catalysis reaction equipment crucibles

IrCl3 H 3 IrCl 6

precursor material for different purposes

9.5 Applications

the heated parts of the substrate surface at temperatures in the range of 250-500°C. CVD noble metal coatings are used in the production of electronic devices. Special advantages are obtained by the application of laser-induced thermal decomposition for the production of patterned films and the deposition of gold films in the circuit repair of thin film circuitry, the interconnection of discrete regions on integrated circuits or modules, and the deposition of bondable gold films on thin film and chip carriers. CVD silver films made from different precursor materials enable the selective metallization of solar cells. Thin silver films limit the diffusion of superconducting YBa 2 Cu 3 O 7 layers and improve the electrical properties (Kodas and Hampden-Smith, 1994).

9.5 Applications 9.5.1 General

The applications of noble metals cover a wide range of different areas. (Table 9-46). Figure 9-35 indicates the quantities required in 1993 for the main application areas. Noble metals are applied in their elementary, unalloyed form in catalysis (platinum, palladium, rhodium), and in electronic, sensor, and optical devices (silver, gold, palladium, platinum). For the majority of applications they must be alloyed for strengthening or to adjust their thermal, physical, or chemical properties according to special requirements. The prices of noble metals are subject to fluctuations, which are determined by the availability on the market from either noble metal production or sales of reserves, and by the demand for applications requiring larger quantities. Examples include

545

Table 9-47. Noble metal demand and price in the western world 1993 (Degussa, 1994; Johnson Matthey, 1994). Metal Au Ag Pt Pd Rh Ru Ir Os

Quantity (10 6 troyoz)

Average price ($/troy oz)

Total (106 $)

72.58 514.09 4.04 4.26 0.36 0.18 0.033 0.005

359.79 4.31 374.06 122.36 1100.01 19.22 102.45 400.15

26113.1 2215.7 1511.2 521.2 396.1 3.46 3.4 2.1

gold for jewelry, silver for photography, platinum and rhodium for automotive catalysts, and palladium for electronic devices. Table 9-47 gives an idea of the commercial importance of the noble metal market with the average prices in 1993, the total demand in the western world, and its monetary value. New applications (e.g., platinum for fuel cells) may change the relative amounts or generate further, individually growing demands. The applications of noble metals can be classified into three main groups: a) Nonindustrial applications, which mainly include the use of gold, silver, platinum, and palladium in jewelry and dentistry. b) Applications as components or connecting materials in devices in the fields of electrical engineering, electronics, and chemical technology, and optical and heat reflection. Gold, silver, platinum, palladium, and ruthenium are predominantly used. c) Applications with combined physical and chemical interactions: Interactions can occur by the exchange of electrons, by charge transfer between adsorbed atoms or molecules and noble metal surfaces, or by intermediate chemical reactions of compounds in solutions. The corresponding

546

9 Noble Metals and Their Alloys (e)

(a)

Chemical 30%

Jewelry 87% Electronic 6% Remainder 4 % Dental 3%

Other 2%

Electrical 68%

Film/photography 40%

(f) Catalytic converter 90% Solder 5%

Jewelry and silverware 26%

Remainder 15%

(0

Chemical 3 Electrical \ Glass 1% Other 4%

Electronics and electrical devices 14%

Catalytic converters 35%

(g)

Other 32%

Chemical 5% Electrical 4 % Glass 3%

Other 4 % Petroleum 2%

Investment 8%

Crucibles 6% Chemical 62%

Jewelry 40%

(d) Dental 29%

Chemical 4 % Catalytic converters 15%

Other 1 % Jewelry 5%

Electrical 46%

Figure 9-35. Main applicational areas of the noble metals, (a) gold (Degussa, 1994), (b) silver (Degussa, 1994), (c) platinum, (d) palladium, (e) ruthenium, (f) rhodium and (g) iridium (Johnson Matthey, 1994).

9.5 Applications

applications are heterogeneous and homogeneous catalysis, and electrode processes in electrochemisty, batteries, and fuel cells. Predominantly used noble metals are the platinum group metals and in special cases (oxidation reactions) silver. Important catalytic functions are performed in addition by complex noble metal compounds in the homogeneous catalysis of numerous organic chemical processes. 9.5.2 Main Application Areas 9.5.2.1 Non-industrial Applications Jewelry Gold and silver have played an important role since ancient times as jewelry metals, silverware, medals, and coins. This trend has continued up to now with a high level of demand, the quantity and value of noble metals for these applications representing an important commercial factor in the noble metal market. In addition to this, platinum and palladium are also becoming increasingly important in jewelry. Noble metals are used in jewelry in the form of a variety of different alloys. The choice of alloy is determined both by aesthetic aspects and technological requirements concerning the working properties for industrial production and for wear and corrosion resistance. Gold jewelry alloys are usually classified as colored or white-gold alloys (Table 948) (Drost and Hausselt, 1992). Colored gold alloys are based on the ternary AuAg-Cu system (Fig. 9-36). This alloy system offers a variety of different colors and the possibility to adjust the mechanical properties by the composition according to requirements. The colors vary depending on the gold content and the relative proportions of silver and copper at a con-

547

Table 9-48. Main components (wt.-%) of gold-based jewelry alloys. Colored gold Fineness 750 585 375 (333) White gold Fineness 750 585 375 (333)

Ag

Cu

0-20 5-35 5-15 5-40

5-25 5-35 45-50 25-60

Ag

Cu

Pd(Ni)

0-10 0-25 0-35 0-35

0-10 5-30 5-50 10-50

10-20 5-20 5-20 5-25

Zinc max. 20%, tin, indium, gallium each max. 4%.

stant gold content (e.g., 18 Ct or 14 Ct) from gold-yellow to copper-red and silverwhite. Copper-rich gold alloys at a level of 14 Ct or even lower are usually alloyed with zinc (max. 15%) to change their red color to a yellow hue. The colors of gold alloys can be determined by spectrographic methods. For practical purposes, typical colors of special alloys are described by special parameters for "tone" (T), "saturation" (S), and "darkness" (D). Methods for their determination by spectroscopy and values for different color grades are standardized (DIN 5033, 6164, 8238). The mechanical properties of ternary Au-Ag-Cu alloys are determined by the wide miscibility gap in this alloy system. Hardening occurs by decomposition of the homogeneous Au-Ag-Cu phase below the critical temperature into a copper-rich (Cu-Au) and a silver rich (Ag-Au) phase. In high-copper-containing 18 Ct gold alloys the formation of the ordered Au-Cu phase contributes to the hardening. For working purposes, the alloys can be obtained in softer conditions by quenching from temperatures above the phase transi-

548

9 Noble Metals and Their Alloys

Au

Gold-yellow

0 100

Green-yellow

^^

/

Ag<

\Vi8Cf &YellovA

// egreenish] ale jYellowish A yellow 1

r

T\

^

y

8/ 0 /

100

Y

/"A

^

/ /

20/ /

Whitish

!

20

J

\^o

\ -8"

\ / / Reddish

Red

\ Figure 9-36. Color ranges in the Au-Ag-Cu system (Renner, 1990, 4).

I

40 60 Copper (wt.%) — —

100Cy

80

Graduation of color of gold-, silver-, copper alloys

tion temperature. They can be rehardened by tempering at temperatures up to slightly below the phase transition point. The binary alloy of 99 % Au and 1 % Ti shows the bright color of pure gold. With its high caratage (23.76 Ct), it is primarily conceived for the top segment of the jewelry market. The hardness, already 70 HV in the soft condition (compared to 40 HV for fine gold), can be raised by work hardening up to 120 HV and by subsequent age hardening at 500 °C up to 170 HV, leading to a fine distribution of the intermetallic TiAu4 phase (Gafner, 1989). White gold alloys contain either nickel or palladium as a bleaching agent. They are typically used with gold contents of 14 Ct and 18 Ct. Binary gold-nickel alloys reach considerable hardness by solid-state separation below 821 °C into gold-rich and nickel-rich phases. Copper and zinc additions improve the formability of nickelwhite gold alloys. Color corrections are made by plating with rhodium. Nickel-

white gold may cause allergic reactions. The trend is therefore towards palladiumwhite gold alloys. These are advantageous in their malleability, ductility, and corrosion resistance. As palladium raises the melting temperature in gold alloys, they are alloyed with appropriate amounts of silver, copper, nickel, and zinc to lower the melting range and increase the hardness. Silver alloys for jewelry and silverware contain up to 20% Cu. Alloys with a fineness of 925/1000 (sterling silver), 835/1000, and 800/1000 are the main types. Their low hardness in the soft condition (HV 70-85) and the high ductility of sterling silver makes them best suited for sledge, bend, and filigree work. Hardening is performed by annealing at 290 °C, causing the segregation of copper atoms in the grain boundaries of the silver-rich phase. The harder 800/1000 alloy is preferred for the production of cast items with a higher scratch resistance. For enamel work the composition 970/1000 is used to avoid lo-

9.5 Applications

cal melting at heating temperatures exceeding 780 °C. Silver jewelry has to be protected against corrosion by coating with clear lacquers or by galvanic plating with rhodium. Platinum and palladium alloys have gained growing importance in jewelry, and also for high quality watch cases and spring parts. The minimum fineness of the alloys is 850/1000 when alloyed with copper and cobalt as hardeners, and iridium, palladium, and ruthenium as grain refiners. Most of the alloys contain 95% platinum, hallmarked PLAT 950. Special alloys are in use for encasing gemstones (Pt-Co/ 950), for watch cases (Pt-W/950), and for spring-parts (Pt-Ir/900). Jewelry Solders Individual chain links, settings for stones, and hollow ware parts are soldered with solder alloys, which must correspond in caratage, color, and mechanical properties to the materials to be joined, but have melting temperatures sufficiently below these to avoid melting the noble metal pieces. The solders and brazing alloys used are principally of the same or similar composition. For consecutive solder operations, alloys with different melting ranges (first, second, and third, solder) are applied. Joining at considerably lower temperatures is possible by diffusion soldering. This process is a hybrid of diffusion bonding and soldering. The principle is to apply a liquid solders, which solidifies at elevated temperatures by isothermal reaction with the substrates to be jointed, so forming solid phases with higher melting temperatures. Tin can be used to make diffusion soldered joints between components of pure gold (see also Sec. 9.5.2.2) (Humpston et al., 1993).

549

Silverware Silverware includes a variety of practical and ornamental items. For general purposes they are usually classified as flatware (tableware such as spoons, forks, and general cutlery) and hollow ware (bowls, water pitchers, coffee pots, etc.). The alloys usually employed are Ag-Cu grades of 958 (Britannia silver), 925 (sterling silver), 835, and 800 silver fineness. On ancient pieces, the hallmark designations are: "10-lot" = 625 fineness, "12-lot" = 750 fineness ("Berlin silver", mark: scepter), "U'/i-lot" = 781.25 fineness ("Frankfurt silver, mark: eagle), "DVi-lot" = 834.75 fineness ("Zurich silver, mark: Z), "14-lot" = 875 fineness, and "14-lot, 4 grains" = 888 fineness ("Venetian silver"). Special techniques are applied for finishing cast and deep-drawn pieces, e.g., embossing, punching, chasing, and engraving ("guillochising"). For complex shapes the lost wax method is also applied. Because of the relative softness of high-silver-containing alloys, the majority of tableware is made from silver-plated, harder base materials. Special galvanic silver-plating techniques and bath compositions have been developed for these applications (Drost and Blass, 1995). Coins, medals, and bars Although special high gold- and silvercontaining coins are still declared and issued as legal currency in some countries, noble metal coins have practically no or only minor importance in monetary use. They are, like medals and bars, primarily collector and investment items. In silvercontaining money, the silver content has

550

9 Noble Metals and Their Alloys

Table 9-49. Selected gold coins: (a) types of NM coins, (b) silver coins, (a) (Renner, 1990, 2) Country

Coin

20 mark 10 mark 5 mark

Mintage period

) r German Reich

}

J

J

1 ducat Kruegerranda Maple Leafa American Eaglea Britanniaa Nuggeta Tscherwonez

German Confederation Rep. South Africa Canada U.S.A. Great Britain Australia U.S.S.R.

Also in | , ± and ^ oz gold;

b

> 1871- 1915

up to since since since since since since

1871 1967 1979 1986 1987 1987 1975

Fineness

Carat

Gross weight (g)

Gold content

900 900 900

21.6 21.6 21.6

7.964 3.982 1.991

7.168 g 3.584 g 1.792 g

986.1 916.6 999.9 916.6 916.6 999.9 900

23.7 22 24 22 22 24 21.6

3.490 33.931 31.103 33.931 33.931 31.103 8.60

3.441 g 1 oz b 1 oz b lozb 1 oz b 1 oz b 7.74 g

1 oz = 28.35 g.

(b) (Renner, 1993, 2) Country

Year Issue

United States

Face value

Germany

ca. ca. ca. ca. ca. ca.

1895 1895 1895 1895 1895 1895

$1 $1/2 $1/4 $1/5 $ 1/10 $ 1/20 $1/2 $1/2

900 900 900 900 900 900 900 400

24.05 11.25 5.62 4.50 2.25 1.13 11.25 4.60

ca. ca. ca. ca. ca.

1895 1895 1895 1895 1895

sfr. sfr. sfr. sfr. sfr.

900 835 835 835 835

22.50 8.35 4.175 2.083 12.525

900 900 900 900 900 900 500 500 500 900 625 625 625 625

25 10 5 2.5 1.0

1873 1873 1873 1873 1873 ca. 1880 1924 1924 1924 1936 1936 1951-1974 1952-1979 1972

Silver per coin (g)

Circulation

1964 1965 Switzerland

Fineness

5 2 1 0.5 5

M5 M2 M1 M0.5 M0.2 M 5/2/1/0.5/0.2 RM 5 RM2 RM 1 RM 5 RM 2 DM 5 commemorative coins DM 5 Olympic Games DM 10

Total silver content of coins issued (t)

6000 10 5 2.5 12.5 7.00 7.00 9.6875

2450 500 1850 1100 870

9.5 Applications

gradually been reduced, and silver-clad coins have also been introduced; the silver content of these coins varies to a considerable extent. Table 9-49 gives some old and some still currently issued coins based on gold and silver. Dentistry The use of gold for repair and prothetic parts in dentistry goes back as far as « 500 B.C. Gold, silver, palladium, and platinum are the main constituents of today's dental alloys in both conservative (for the filling of cavities and inlays) and prosthetic (crowns, bridges, metallic basis for ceramic blends, etc.) dentistry. Dental alloys are available in different grades of mechanical strength and melting ranges (Kempf and Hausselt, 1993). Material for dentistry repair and replacement has to be corrosion resistant, sufficiently wear resistant, heat conducting, and may show no or only minor shrinkage after hardening. Besides ceramics and composites, noble metal alloys form the major part of dental materials. In conservative dentistry more than 50 % of cavity fillings are made of silver amalgams. They are formed from pasty mixtures of Ag-Cu-Sn alloy powders and mercury, which harden at room temperature forming a complex matrix of silver amalgam and base-metal compounds. Typical compositions are in the range of 40 wt.% silver, 32 wt.% tin, 30 wt.% copper, 2 wt.% zinc, and ^ 3 wt.% mercury. Special gold-flake preparations can be used to fill small cavities by means of ultrasonic wave densification. In prothetic dentistry (crown and bridge forming), in addition to gold alloys, silverbased and palladium-based alloys are also in use. Alloys provided for successive surface sealing with ceramic veneers (PFM

551

alloys: porcelain fused to metal alloys) are conceived for their adherence on porcelain and have their coefficient of thermal expansion adjusted to that of the ceramic. Basic types of noble metal alloy components for prothetic dental alloys are given in Table 9-50. The available hardness of dental alloys ranges from 50 to 275 HV. Alloy additions of silver, palladium, platinum, copper, tin, indium, zinc, and gallium affect the hardness, melting range, and adhesion to ceramics. Iridium and rhodium (to a lower extent also rhenium, tantalum, iron, and ruthenium) are used for grain refining of gold-based alloys, ruthenium and germanium for grain refining of palladium-based alloys. The fine grain structure improves the mechanical properties and also, because segregation effects are minimized, the corrosion resistance. Hardening is achieved in high-gold-containing Ag-AgCu alloys by precipitation hardening. The hardness is adjusted by varying the silver and copper contents, causing changes in the amounts of precipitated Au-Cu and Ag-Cu phases, so making it possible to adjust the hardness according to the required properties. The melting temperatures of Ag-Au-Cu alloys range between ^800 and 1300°C. Table 9-50. Basic compositions of dental alloys. Noble metal base

most common alloying elements

Crown and bridge alloys Au Au-Ag Ag-Pd

Ag, Cu, Pt, Zn Pd, Cu, Zn, In Cu, In, Au, Zn

Porcelain fused to metal alloys Au Au-Pd Pd Pd-Ag

Pt, Pd, In, Sn Sn, In, Ga, Ag Cu, Ga, Sn, In Sn, In, Zn

552

9 Noble Metals and Their Alloys

The upper value of practical usage is limited by the melting capability of the operating equipment. Alloys provided for surface sealing with ceramic veneers must have solidus temperatures above 1000°C to withstand the sintering temperature of the ceramic blend. Alloying with platinum (minimum 6 wt.%) and palladium increases the hardness and the melting temperature. The adherence of the ceramic is aided by alloying with base metals which form oxides on the boundary metal/ceramic during bonding. The thermal expansion of the alloy must surpass that of the ceramic to a small extent. This generates a compressive stress on the ceramic after cooling, which enables the combination to withstand repeated temperature changes without crack formation. To maintain sufficient corrosion resistance, the noble metal content of each separate phase must exceed 50%. Palladium-based alloys (palladium content from ^60-80 wt.%) contain copper and gallium as hardening elements, and tin to lower the melting temperature. The best corrosion resistance is obtained with a single-phase structure in the composition range of Cu 4-6 wt.%, Ga 5-7 wt.%, and Sn 6 wt.%. Palladium-silver alloys contain ^ 3 0 wt.% Ag. Additions of tin, indium, and zinc harden by solid solution or precipitation, and lower the melting range. Prothetic dental pieces are manufactured by casting or by powder-metallurgical methods. Joining is done by the techniques of soldering, microplasma welding, or laser welding. Soldering is the most commonly used process. The solder alloys are of similar compositions to the dental alloys. They must correspond to the base materials in their corrosion behavior and thermal expansion. Solders for crowns and bridges melt in the temperature range of «760-840°C. They are based on the Au-

Ag-Cu system with additions of zinc and tin. Melting is done by gas torch. Solders for PFM alloys are used before porcelain veneering. The melting temperatures and softening ranges must be sufficiently high so that subsequent baking does not influence the mechanical stability of the solder joint. The solders are based on the Au-Ag system with reduced palladium and platinum contents. Color lightening is effected by palladium additions, and the melting temperatures are adjusted by additions of tin, indium, or zinc. Biomedical Uses Numerous applications exist both for noble metals and noble metal compounds in biomedical usage. As nearly all noble metal compounds are highly toxic, their application is limited to special diseases and requires additional care and supporting treatment with supplementary medicines. Silver compounds are only occasionally used in medical therapy, while previously they were used to treat epilepsy and ulcers. Silver complexes are used to disinfect the pharynx because these complexes are readily absorbed by the mucus membranes. The local etching effect of silver nitrate is used to destroy exuberance in tissue. The compound sulfadiazine silver (i), [AgSD = 2-sulfanilamido pyrimidine silver (i)] stops the formation of bacteria after severe burning. Silver and silver alloys are used in brain surgery as a substitute for bone. Silver wire and silver gauzes are implanted. Gold was the first noble metal to be used in chemotherapy. Paracelsus was the first scientist to recognize the therapeutical effect of organic gold compounds for rheumatic ailments and arthritis. Various complex organic gold compounds are still in use to treat rheumatoid arthritis.

9.5 Applications

Platinum finds many important applications in medical care (Kidani, 1991). The catalytic properties of platinum and the PGMs are exploited in the synthesis of special drugs, including pain killers, antibiotics, or asthma medicines, and for the synthesis of organic stereo-enantiomer compounds. A complex organic rhodium compound is used as a catalyst for the production of the amino acid "L-dopa", which has become important in the treatment of Parkinson's disease. The stereo-isomer dichloro-diammino Pt n -complex compound ("cw-platinum") os-[PtCl 2 (NH 3 ) 2 ], NH 3 .

Cl Pt1!

NH,

553

cancer therapy. The radiopacity of the platinum shields the healthy tissues from radiation, while the tumor is irradiated by the exposed iridium tip. Platinum coils made from thin wire are applied for the treatment of aneurysms in the brain, where traditional recourse of surgery is most difficult (Coombes, 1993). Another area of medical application for noble metals is acupuncture needles, made from high carat gold and silver alloys. 9.5.2.2 Components and Connecting Materials Electrical Engineering General

*C1

[the racemate is "Peyrone's chloride" (Peyrone, 1845)], which was first isolated by Werner (1890), was recognized as a powerful agent against cancer by Rosenberg etal. (1965, 1969). The anti-tumor effect can be explained as a chemical bonding of Pt11 onto two nitrogen atoms, each belonging to a nucleo-base of the two molecular strains of the DNA double helix. Guanine, one of the four nucleo-bases of DNA combines readily with the Pt n -complex. Succeeding products with lower toxicity ("Carboplatin", "Paraplatin"), to be applied by infusion, are in use, and newly developed derivatives, to be taken orally, are in the test stage. Because of their excellent corrosion resistance combined with electrical conductivity, platinum and platinum-iridium alloys find wide use in medical devices, such as electrodes for pacemakers, catheters, marker bands, guide wires, and as temporary or permanent implants. Irradiated iridium wire sheathed in platinum can be implanted into the body to deliver controlled doses of radiation for

Noble metals are essential components of materials for electrical contacts, connectors, conductors, resistors, and fuse elements. Electrical contacts serve to close circuits, to take over the current conduction, and to reopen closed circuits. The contact material has to be stable under working conditions and must not change its characteristics over a large number of switching cycles. The absence of tarnish layers is of primary importance for switching contacts working at low voltage and low current with low contact loads. Requirements for heavy current contacts are thermal erosion resistance, arc extinguishing performance, low tendency for welding, and low material migration. Applications are covered by a large variety of material compositions for high and low current technology. Because of the low strength of pure noble metals and their tendency to weld under mechanical and thermal loads, in most cases alloys or noble metal bearing composite materials are used (Schroeder, 1995).

554

9 Noble Metals and Their Alloys

Make-Break (Switching) Contacts Silver, silver alloys, and silver-bearing composites make up the largest amount of contact materials and are especially dominant in high current technology. Typical compositions are listed in Table 9-51. They cover a wide range of properties, which are related to the special requirements of the device, for example, surface welding during make, resistance of close contacts, arc migration, and arc erosion during break. Fine silver is used as a galvanic coat in thicknesses between 2 and 50 \xm to lower the contact resistance and to improve the Table 9-51. Silver-bearing composite contact materials. Type

Alloys

Composition Hardness (HV) (at.%)

Electrical conductivity (m/Q-mm 2 )

100 120

58 52

150

49

90

54

Ag-Ni(40)

115

37

Ag-CdO(lO)

80

48

AgNi(0.15) AgCu(3) 1 AgCu(20)

Composites with: Metals Ag-Ni(lO) i

Oxides

i

i Ag-CdO(15) Ag-ZnO(88) Ag-SnO2(8) Ag-SnO2(12) Carbon Refractory metals/ compounds

115 95 92

45.5 49 51

100

42

Ag-C(2) Ag-C(5)

40 40

48 43.5

Ag-W(20)

240

26-28

80 130

42 24-30

470

-

i Ag-W(80) Ag-WC(40) i

I Ag-WC(80) a

Composite made by infiltration of liquid silver into a tungsten skeleton.

solderability of base-metal contact pieces. The hardness is between 60 and 160 HV; the electrical conductivity amounts to 5080% of that of the cast metal. The Ag-Ni and Ag-Cu alloys (fine grain and hard silver grades) are soldered or clad on copper alloy backings for switches in DC and AC devices working at currents of up to 10 A. They have a higher wear resistance and can easily be soldered with common silver solders. Ag-Ni composites are used for AC and DC devices which work in the low voltage range and at currents below 25 A. They can also be soldered or welded without extra surface-plating. Silver-metal oxide (CdO, SnO2, ZnO) composites have excellent resistance against welding during make and in the closed position under current. They cover the range of 50-5000 A in AC switches especially with high current peaks on make, e.g., low voltage motor switches, power circuit breakers, and relays in automobiles. The composites are produced either by mixing the metal and oxide powders, densification, sintering, and extrusion, or by internal oxidation of the corresponding silver (cadmium, tin, zinc) alloys in air at temperatures of 700-850 °C. Internal oxidation of the silver-tin alloy requires pre-alloying with indium (1.5-4 wt.%) to crack the dense SnO2 surface layers which form during oxidation and prevent the inward diffusion of oxygen. Discrete single-contact pieces are made from internally oxidized silver-metal powders by pressing these into the final shape with subsequent sintering. Silver-graphite composites with 2-5 wt.% C are the preferred material for sliding contacts. They offer high protection against contact welding under short circuit currents. Extruded rods show similarly to silver-nickel composites, a fiber-like orien-

9.5 Applications

tation of the carbon in the direction of the rod axis. Contact pieces are cut along the cross section, so that the carbon fibers are perpendicular to the contact surface. AgMo, Ag-W, and Ag-WC composites are used for power switches which work at very high temperatues. Nickel additions serve as the wetting and bonding agents. Ag-W composites with tungsten contents above 60 % are made by infiltrating liquid silver into a sintered, porous skeleton of tungsten. Infiltration takes place by capillary action. Because of their higher corrosion resistance, gold and platinum metals are preferentially used in components for communication and information equipment where even thin oxide layers can cause interruptions or failures in signal transfer. For low power applications, gold- and platinumbased contacts are used. Materials for medium power (relays in telecommunication or automotive applications) are Pd-Ag and Pd-Cu based contact alloys. In Pd-Cu40 an ordered structure with lower electrical resistivity can be formed by thermal treatment. The make-break properties of gold are improved by alloying with platinum, silver, nickel, or copper. Due to their catalytic properties, palladium and platinum tend to form organic polymer compounds ("brown powder effect") on the contact surface from hydrocarbon compounds present in the atmosphere, which disrupt the contact behavior. These reactions can be reduced by alloying with silver, iridium, tungsten, or nickel. Pt-Ir (with iridium contents of 10-20 wt.%) is especially stable against surface burning, and Pt-W5 and Pt-Ni8.5 are very stable against spark erosion Stationary Contacts Stationary contacts are used in connectors, plugs, terminals, and sockets. They

555

contain noble metals and their alloys as surface layers on metallic supports. They are clad by rolling, pressing, or galvanic (e.g., selective) coating. Small additions of cobalt or nickel in the galvanic baths increase the hardness and wear resistance of gold deposits. Nickel underlayers favor the formation of dense, pore-free, thin gold coatings. Sliding Contacts Sliding contacts are used for collectors in motors, rotary switches, brushes, and resistance wires for potentiometers, commutators, and slip rings. They are subject to larger wear due to sliding under electrical current. Common material forms are clad bands or profiled strips (overlay, inlay, toplay, or similar design). Ternary and quaternary alloys of Pd-Au-Ag-Pt and PdAg-Cu, and for hardest wear resistance alloys of Pd-Co-W and Pd-Co-Cu, are clad. Fuses Fuses have to protect electric lines against current overload. Fuse elements are designed to melt, breaking the current flow if overload or short circuit develop. According to the different requirements, a large variety of different types and structures are used. Applications include voltage ranges from several volts to several tens of thousands of volts, and currents from 2 mA to several thousand amps in rated current. Silver is the dominant noble metal, and is applied in different shapes, either in pure or alloyed form (for example, with 50 wt.% copper) or clad on copper with silver layers corresponding to a silver content usually between 0.6 and 20 wt.%. Special constructions in combination with soft solder coatings effect retarded or fast melting. The simplest shapes consist of wires with a reduced diameter

556

9 Noble Metals and Their Alloys

over a defined length, which melt under overload. Conductor and Resistor Materials Silver is used as a conductor material on copper or aluminum supports for current leading busbars. In AC technology at medium and high frequencies, based on the skin effect, silver coatings of 1-5 jam thickness on copper conductors are the current-bearing media. Application areas are UHF (ultrahigh frequency) television and radar technology, coaxial cables, and space technology. Surface protection is provided by thin coatings of either gold, palladium, or rhodium. Resistor-heated furnaces working at temperatures > 1200 °C are equipped with heaters made from platinum wires (up to 1600°C), Pt-Rh40 (up to 1700°C), or rhodium bands (up to 1800 °C). These furnaces can be used in an air atmosphere. Rhodium and iridium are suitable for susceptors for medium and high frequency fields for inductive heating in air, which is used for the induction melting of glasses or ceramics in crucibles made of iridium, which have been protected against oxidation by coating with rhodium. Au-Cr2 serves as a resistor material in corrosive environments with a stable TCR between -20 and +40°C. Wires of multicomponent alloys based on gold or palladium serve as potentiometer windings. Their specific resistance values range from 15 to 150 JIG cm with a TCR extending from + 1 0 " 4 to - 1 0 " 5 (per °C). Sliding contacts are made from Pd-Ag or Ag-Au-Cu alloys. Electronic Applications General

Noble metals find extensive applications in electronic devices and circuits as functional components:

• conductors, electrodes, resistors, or contact plates, • surface layers for electrical contacts and corrosion protection, • constituents of soldering and bonding materials for electrical and mechanical connections. Electronic applications comprise vacuum electron tubes, active and passive semiconductor devices, electronic circuits, components, and the materials needed for their production. Table 9-52 gives a general overview of the applications of noble metals in devices, components, and materials. Silver is, due to its high electrical conductivity, the preferred conductor and electrode material. Large quantities are used in oxide-ceramic based devices such as electrode layers and contact metallizing, and as a conductor material in thick film and thin film hybrid circuits. It is a constituent of many vacuum brazing alloys for the manufacture of electron tubes and ceramic and metal packages. Silver additions to soft solders based on lead and tin increase the strength, improve the wetting and flow properties, and reduce leaching of the silver surfaces in solder processes. Gold is primarily used if oxide-free surfaces are indispensable, enabling perfect solder, bond, and pressure contacts. Partly due to its high intrinsic cost, it is mostly employed in the form of thin layers, which are preferably deposited by electroplating. Different galvanic bath compositions are used for high purity gold coatings for bonding, and hard gold coatings for connector applications (Christie and Cameron, 1994). Gold and gold-silicon alloys are used to connect the semiconductor crystals with the base contact faces ("dieattach") on lead frames, packages, or hybrid circuits. The most common process

Table 9-52. Applications of NMs in electronics. Form

Element Bulk

Bimetal

Surface coat electroplated

sputtered or evaporated

pastes, lacquers

solder connection

Ru

reed contacts

Ru-0 for cathodes in special tubes

Mo/Ru hightemperature solders for electron tubes

Rh

slide contacts, coating of Cu discs for thyristors, surface hardening of Ag coatings on pressure contacts plug-in contacts

hard contact layers for pressure contacts on Ag surfaces, corrosion protection

thick film circuits (resistor pastes for hybrid circuits) metallizing pastes for sensor surfaces

Ag-Cu-Pd solders for electron tubes package joining

printed circuits boards (PCB) for high frequencies, system carrier, Mo and W base-plates, lead frames

contact surfaces on Si diodes lead frames

lambda probe, conductor and electrode paste for thick film circuits, multilayer capacitors conductor pastes for thick film circuits, potentiometer, capacitors, contact foils, screening, glass solders

Pd

bonding wires for transistors, diodes

plug-in contacts

Ag

pressure-contact plates for power semiconductor devices

plug-in contacts, pressure-contact and soldered base plates

Os

melting crucibles for laser crystals (YIG, YAGa) melting crucibles for laser cystals (YIG, YAGa) test resistors

Ir Pt

Au

bonding wires for integrated circuits transistors, diodes

multilayer capacitors

tubes, package joining components, soft solders in semiconductor devices, die-attach

CO

plug-in contacts

PCBs, plug-in contacts, solderable and bond surfaces, switches, corrosion protection of semiconductor device packages

YIG = Yttrium iron garnet, YAG = Yttrium aluminum garnet.

test resistors, thin film sensors, diffusion barriers, IC technology thin film hybrid circuits, conductor contact, bond surfaces

lambda probe, thick film sensors thick film hybrid circuits, solderable areas for contacts, dieattach, IC technology

special solders for dieattach, package soldering, and connections in semiconductor power devices Ol Ol 1

558

9 Noble Metals and Their Alloys

for connecting the semiconductor contact pads with the inner leads of the device is "bonding" with thin gold wires ("bond wire"). Gold-containing brazing and soldering alloys are used for the joining of packaging and assembling components. Special packages (flat-pack) are hermetically sealed with gold-plated lids, using solder preforms of the alloy Au-Sn20. Palladium, whether pure or in combination with silver, is the most important component in capacitor electrode layers, in silver conductor paths, and in resistor elements in thick film circuits. Platinum finds applications as barrier layers, contact pads, and stabilizing components in silver conductor materials. Ruthenium has become the most important base element of resistor elements. Rhodium serves as as hard, rear-resistant, electrically conductive surface coating for slide and pressure contacts. High purity (standard 99.99 % or, if the material is in contact with the semiconductor crystal, 99.999%) and high accuracy of the alloy composition are required. For ohmic contacts at the semiconductor, surface layers and solder alloys have to be doped (low; defined as alloyed) with elements corresponding to the p (positive) or n (negative) type of semiconductor to minimize contact resistances. Common alloying elements are gallium or indium for p-type alloys, and arsenic or antimony for n-type alloys. Devices Vacuum tubes consist of a high vacuum chamber, usually made from A12O3, with inner electrodes which are vacuum tight, and electrically and mechanically connected by outer leads. Joining is performed with high purity silver solder alloys of low carbon content (<10ppm) and free of

gases or volatile components. Before soldering, the ceramic surfaces are usually metallized by coating with manganesemolybdenum pastes, then fired and galvanically plated with nickel. Titaniumcontaining, "active" silver solder alloys allow direct soldering without preliminary metallizing. All solder processes are performed under a protective gas atmosphere (argon, nitrogen, hydrogen, mixture of nitrogen and hydrogen) or in a vacuum. Active semiconductor devices contain semiconductor material (silicon, germanium, gallium arsenide, ceramic oxides, e.g., BaTiO3). Monocrystalline semiconductor materials (chips) are used for integrated circuits (IC), diodes, transistors, and power devices; polycrystalline and amorphous semiconductor materials are used as constituents of solar cells and of oxide semiconductor devices. Thin films of gold, silver, platinum, and palladium, and the silicides of palladium and platinum in combination with barrier layers of titanium, chromium, and molybdenum serve as wiring and electrode materials for interconnection. Oxide semiconductor materials and passive devices like capacitors, resistors, and potentiometers are electrically and mechanically connected by applying silver-metallizing preparations followed by firing. Important passive devices are the multilayer capacitors and discrete chip resistors made from ruthenium compounds. Multilayer capacitors are built up from alternate layers of ceramic and a conductive electrode material, usually consisting of palladium or Ag-Pd alloys. The layers are terminated at each end of the chip by a silver coating (Fig. 9-37). Production is done by applying ceramic and conductor pastes alternately, laminating together ("wet-wet" process), and firing in one single operation.

9.5 Applications

Figure 9-37. Multilayer capacitor (cross-section) (Demetron).

Circuits

Electronic circuits consist of layer arrays of conductor and resistor paths, loaded with active and passive devices. Besides the single-crystal integrated circuits (ICs), the main types are printed circuits (PCs) in single- and multilayer arrays, and thick film and thin film hybrid circuits. Pure gold and silver layers of 1-2 jim thickness form on PC board conductor

Gold bond pads

559

paths and contact pads for die-attach ("chip on board"), or for connection with bonding wires. Hard gold coatings are applied for connector and slide contacts. Electroless (by chemical reduction) deposited, thin palladium coatings form the first metallization step for metallizing bare polyimide surfaces, followed by chemical and galvanic deposition of thicker copper layers, as applied for the plating of holes to connect electrical functions in multilayer PCs. Thick film circuits (Fig. 9-38) are produced on ceramic substrates, using thickfilm pastes containing silver, gold, palladium, platinum, or ruthenium powders of different compositions with auxiliary organic materials and special grades of glass powders. Conductor and resistor paths are produced by screen printing, drying, and firing. Thin film circuits are designed for low power applications. Conductor and resistor networks are produced by combined sputtering and etching techniques, leading to very fine structures with high packaging density.

Gold bond wires

Edge metallization of multilayer capacitor (Ag, Ag-Pd) Silver adhesive for "die-attach"

Inner metallization of multilayer capacitor (Pd, Pd-Ag)

RuO2 resistor

Conducting paths (Ag, Au, Ag-Pt, Ag-Pd

Figure 9-38. Thick-film circuit (Dorbath et al., 1991, 2).

560

9 Noble Metals and Their Alloys

Components and Materials Silver and gold find extended use as surface layers and in the bulk form in components (lead frames, base contacts, chip carriers, packages) for chip bonding, assembling, electrical and mechanical connection, mechanical support, and protection against corrosion. Materials applied for the production of devices, circuits, connections, and surface layers are solder alloys, sputter targets, evaporation slugs, bonding wires, thick film pastes, and galvanic baths. Silver- and gold-based binary, ternary, and quaternary solder alloys are used for the mechanical and electrical connection of semiconductor metal and ceramic components. Soldering processes are usually done in consecutive steps. The melting ranges of solders used in electronics extend from soft solder ranges (alloys based on lead, tin, or indium, melting between ^ 159 and 400 °C), special eutectic gold alloys (melting between 217 and 356°C), to hard solder Ag-Cu, Ag-Cu-Pd, and Au-Ni alloys with melting temperatures above 600 °C. Soldering processes in electronic device and package production are predominantly carried out without a flux in a protective atmosphere. The solders are applied in the form of wires, foils, solder preforms, (spheres, disks, etc.) or pastes. The assembly of electronic devices is characterized by joining materials with considerably different thermal expansion coefficients (Table 9-53), which can generate considerable mechanical stresses caused either by the thermal stresses during soldering or by thermocycling of the devices during use. The eutectic alloys of gold with germanium, silicon, and tin combine low melting temperatures with high mechanical strength and high corrosion resistance, and are required for die-attach,

Table 9-53. Thermal expansivity of electronic device components. Material

Si Mo W A12C Cu Ag Ail

Pd Pt Pb Sn

Coefficient of thermal expansion

2.3 4.9 4.2 5.5 14.09 17.04 13.2 11.1 8.99 27.08 22.57

for the hermetic sealing of packages, and for soldering laser crystals. Au-Sn alloy connections are also formed by the thermal bonding of tin-plated leads onto gold bumps, placed on the surface of the semiconductor crystal in tape automated bonding (TAB) technology. Au-Si alloys give solder bonds with the highest stability against thermal and mechanical stresses. Au-Gel2 and Au70-In20-Snl0 can be soldered at 400-500 °C. They are used for the soldering of molybdenum supports in package assemblies and leads of copper or of Fe-Co-Ni alloys. Ternary alloys of Au-Ag-Ge, Ag-Cu-Sn, and Ag-Cu-In cover the temperature range of 400-650 °C. Multilayer solder arrays made by cladding on suitable substrates (e.g., Sn/Ag/Sn, In/Ag or Sn/Au/Cu/Au laminates) enable the production of solder connections by liquid diffusion processes at temperatures slightly above the melting temperature of the lowest melting component (Bernstein and Bartholomew, 1965). The ternary alloy Sn-SblO-Ag25 has an extended melting range and higher tensile strength than common soft solders, caused by the presence of low-melting phases and hard intermetallic compounds (Olsen and

9.5 Applications

561

Berg, 1979). It serves for die-attach in devices which have to withstand mechanical and thermal stresses caused by temperature cycling, e.g., in automotive electronics. Sputter targets and evaporation slugs are used for semiconductor metallizing and thin film processes. High purity (99.999%) is required if the sputtered or evaporated deposits come into direct contact with the semiconductor material. Special equipment, designed for single-wafer and multiwafer coating by DC and by magnetron sputtering, has been developed.

(a)

Bonding wires Wire bonding is the most common method of connecting semiconductor crystals with the leads of packages or supporting frames ("inner lead bonding"). About 80 % of all wire bond connections are performed with gold bond wire, ^ 2 0 % are made with aluminum bond wire. To a small extent palladium bonding wires are also used. The yearly demand for gold bond wire worldwide is estimated to be around 12 tons, corresponding to a total number of ^ 1 0 1 2 single bond connections. Bonding is performed by special welding processes ("thermocompression" and "thermosonic" bonding), forming a bond loop between the connection pads. The wire is feed through a heated capillary and the tip of the outcoming wire is melted to a small ball which is pressed onto the semiconductor pad to form the weld contact ("ball" bonding or "nail-head" bonding, Fig. 9-39). The formation of a sufficient, mechanically stable wire loop requires a high yield strength, and controlled values of tensile strength and elongation. According to the Hall-Petch relation = C7O +

K-d-X12

(9-21)

(b) Figure 9-39. Ball-bonding of semiconductor chips (Demetron).

where

562

9 Noble Metals and Their Alloys

fiber texture formed during drawing to a predominantly <100> fiber texture in the direction of the wire axis. If gold bond wire is applied to aluminum-coated pads, special care has to be taken during thermal treatment to avoid formation of the brittle AuAl 2 compound. Gold-based bonding wire made of the alloy AuTil has a higher strength and a higher mechanical stability at elevated temperatures. Strengthening is achieved by thermomechanical treatment causing precipitation of the intermetallic compound TiAu4. Additional hardening occurs by extended heating at elevated temperatures in air, which effects dispersion hardening by the formation of a fine, dispersed titanium oxide phase by internal oxidation. The wire can be made to have comparable electric properties and bonding characteristics by cladding with a surface layer of pure gold (Humpston and Jacobson, 1992). Others Metallizing pastes and lacquers are used for the production of electrode and contact layers of passive devices, single-layer and multi-layer capacitors, and for conductor and resistor layers in thick film hybrid circuits (Dorbath, 1991). Thick film pastes for the production of hybrid circuits are generally applied by screen printing. Conductor pastes are based on silver, silver-palladium, silverplatinum, palladium, platinum, and gold. Resistor layers are printed with paste preparations made from RuO 2 with additions of PbO and Bi 2 O 3 . They are formed by firing complex ruthenium compounds of high thermal stability and very low temperature coefficients of electrical resistivity. The measuring unit is the so-called sheet resistivity, defined as the electrical resistance of a square area, measured be-

tween opposite edges. Exact resistance values are adjusted by applying small cuts by means of laser beams into the resistor layer ("laser trimming") (Starz, 1995). Sensor Devices Sensor devices transform physical parameters, e.g., temperature, pressure, etc. into electrical signals and thereby them make available for special measurement and controlling processes. Devices that determine the temperature by measuring the electrical conductivity are the most important. Widely used noble metals containing sensor devices are thermocouples and resistance thermometers. Thermocouples generate an electromotive force (emf) by temperature-induced electron transfer on the phase boundary of two different metals joined to each other. Platinum and Pt/Rh alloys are best suited for application at high temperatures because of their thermal stability and available purity. Standard combinations are: 1. 2. 3. 4.

" S " PtRhlO-Pt up to 1600°C, "R" PtRhl3-Pt, " B " PtRh30-PtRh6 up to 1800°C, and IrRh40-Irupto2000°C.

(where S, R, and B correspond to the international standard IEC 584.) The measurement of very low temperatures can be carried out with the systems: 5. AuFe (0.03 at.% Fe)-chromel (from 4.2 K to 273 K), 6. AuCo (2.11 at.% Co)-AuAg (0.37 at.% Ag) or Cu (from -240-0°C), and 7. AuFe (0.02 at.% Fe)-Cu (from -270 to -230 °C). Thermocouples of these systems must be individually calibrated because of the difficulty of reproducing the exact composition as a result of small additions of base metals.

563

9.5 Applications

Thermocouples for heavily corrosive atmospheres are: 8. AuPd40-PdPtl4Au3, corresponding to NiCr-Ni, 9. AuPd35-PtPdl2.5 corresponding to NiCr-Ni, and 10. AuPd46-PtM0 corresponding to Feconstantan. Table 9-54 gives basic values of the thermoelectric voltage of some common thermocouples in millivolts. Figure 9-40 shows the application temperature ranges for different thermocouples. Resistance thermometers are based on the temperature dependence of the electrical resistivity. The preferred noble metal is platinum. The temperature coefficient a of the electrical resistance of platinum in a high-purity condition is 3.927 x 10~3/K (DIN IEC 751). Temperature-sensitive devices are constructed with platinum wires of a definite length and cross section or as thin platinum layers on ceramic substrates. Wire-type resistance thermometers are used for industrial temperature measurement and in special constructions as calibration units. They work in the resistivity range between ^25 and 100 fi, but can only be used up to temperatures of

2003

-j 1500 LI

5iooo EJ

:x

E 500

0 -273 1 2

8

9

3

4

5

6

Figure 9-40. Application temperature ranges for different thermocouples: 1 Cu-Au-Fe 0.02 at.%, 2 CuAu-Co 2.1 at.%, 3 Pallador 1,4 Pallador II and Platinel, 5 Pt-RhlO-Pt (DIN 43710), 6 Pt~Rhl8, 7 RhIr60-Ir, 8 Pt-toughened glass-WTh (DIN 43760), and 9 Pt^ceramic-WTh (DIN 43760) (Degussa, 1967).

^600°C. This temperature limit is due to the mechanical weakness of the very thin wires that are needed for sufficiently high resistance values at higher temperatures. Basic values (i.e., required values for the R-T relation) for platinum resistance thermometers are specified in international standards (DIN IEC 751). For temperature measurements at temperatures below « 20 K, the TCR of pure metals becomes

Table 9-54. Basic values of the thermoelectric voltage of common thermocouples3 (DEGUSSA, 1967). 71 (°C) 0

71 (K)

T2 (°C)

Pt-RhlO/Pt

Pt-Rh20/Pt-Rh5

Rh-Ir60

Pt-elc

Pd-ord

100 500 1000

0.643 4.221 9.570

0-074 1.447 4.921

0.371 2.562 5.495

3.31 20.20 41.65

4.6 27.9 59.6

T2(K)

Au-Co2.1/Cu b

Au-Fe0.02/Cu b

10 20 40

0.044 0.173 0.590

0.093 0.208 0.423

4.2

InmV;

b

at.%;

c

Pt-el = Platinel, Pd83Ptl4Au3/AuPd35;

d

Pd-or = Pallador, Ptlr 10/AuPd40.

564

9 Noble Metals and Their Alloys

too low compared to the absolute value of the resistance. At these temperatures, alloys of rhodium with 0.5 at.% iron or gold with 1.15 at.% manganese are used. Their special properties are based on the Kondo effect (Kondo, 1964). The small amounts of iron or manganese cause, as a result of magnetic effects at very low temperatures, increases in the TCR with decreasing temperature, so that the values of the TCR become measurable again. Film-type resistance thermometers are manufactured by thin film technology. Thin platinum or iridium deposits of 1-2 \xm sputtered on the surface of flat ceramic substrates. The deposited film is transformed to a meandrous resistor path by laser-beam cutting or ionic etching and in the same way adjusted to the required resistance value. The soft metallic surface is protected by coating with a thin layer of glass-ceramic. Path widths are around 20 |im (Fig. 9-41). Iridium layers match the ceramic better in their thermal expansion and are therefore especially suited for low temperature measurements down to -200°C (Diehl et al, 1991, Diehl, 1995). Platinum resistors are also produced in thick film technology by screen printing and firing. The resistor pastes contain 90 wt.% platinum powder together with ethylcellulose and glass powder. Conductor widths are around 400 jum. Platinum contact electrodes are constituents of sensor devices made from semiconducting materials (e.g., SnO2, ZnO, piezo-ceramic, NTC [negative temperature coefficient] resistors) or special ceramics. They serve for the determination of gas concentration, humidity, or pressure. Oxygen-determining sensors ("lambda probes") use the high electrical conductivity imparted to solid electrolytes (e.g., ZrO2) by the presence of oxygen ions,

compared to their normal conductivity due to electrons. If this ionic conductor acts as a separator between two gas spaces, and if it is provided with platinum electrodes on both sides, a difference in the oxygen partial pressure causes the formation of an electromotive force ("Nernst emf") between the electrodes. These probes are used to control the fuel-air mixture in conjunction with automobile exhaust-gas catalysts in Otto engines. Noble metal catalysts are also used in sensors for detecting combustible gases in air or explosive gas mixtures. Heat evolution is normally detected by sensors known as pellistors, which consist of a platinum wire placed in a small sphere of A12O3 (diameter: 1.5 mm). The sphere is saturated with a noble metal catalyst. If a chemical reaction takes place there, a temperature increase occurs that can be measured by the change of the electrical resistance of the platinum wire. Pellistors find applications in explosion protection equipment. Magnetic Storage Thin film cobalt-platinum alloys and thin film cobalt-platinum multilayers play an important role as magnetic media on hard disks in magnetic and magneto-optical storage and recording devices. They are deposited on substrates made of aluminum-silver alloy or glass by electronbeam evaporation or sputtering, forming magnetic anisotropic structures. The information input signal is transferred by means of a read-write head into the magnetic layer of the disk, forming magnetization patterns. The head is placed a narrow distance (100 nm) from the disk surface. Grain morphology, grain size, crystallographic orientation, and the structure of grain boundaries influence the magnetic and recording properties (Gau, 1990).

9.5 Applications

565

(a)

, Pt

• Ceramic

(b)

In magnetic layers for so-called "longitudinal recording", magnetic domains are oriented parallel to the film plane. Coatings consist of cobalt-platinum-chromium alloys in varying compositions (Pt: 6-12 wt.%, Cr: 17-18 wt.%), or cobalt-nickelplatinum and cobalt-platinum-alloys, deposited on substrates of nickel-phosphorus/aluminum and chromium. The coercivity is influenced by the type and structure of the surface material of the substrate. It increases with the grain size of the

Figure 9-41. (a) Platinum layer resistor element (dim. 2 x 2.3 mm) with welded connecting wires, layer resistance adjusted by laser cuts, (b) Enlarged view of platinum layer paths on the ceramic substrate. Path width «15pm(Diehletal., 1991, 2).

chromium underlayer and with increasing in-plane c-axis orientation of the Ni-P underlayers. For optimum noise reduction, the magnetic grains must be morphologically decoupled to reduce the spatial extent of the magnetic clusters by reducing the intergranular exchange coupling (Mahvan etal., 1990; Yogi etal., 1990). In magneto-optical devices, the magnetic domains are oriented perpendicular to the film plane (vertical recording), which enables much higher information storage

566

9 Noble Metals and Their Alloys

densities. The processes to write and to erase are performed by local heating of the magnetic domains above the Curie temperature by means of a laser beam at high power, changing the direction of magnetization by an external bias field. Reading processes use a laser beam at lower power, without changing the magnetization. The different magnetic orientations are read by means of the polar Kerr effect, in which the plane of polarization of the reflected laser beam is rotated to the left or to the right, depending on the magnetic orientation of the domain from which it is reflected (Carcia and Zeper, 1990). Magneto-optical devices were originally developed based on rare-earth/transitionmetal alloys (Tb-Fe-Co or Gd-Tb-Co). These alloys give sufficient performance at a higher wavelength of the reading laser, but the effect of the Kerr rotation diminishs if the wavelength changes from infrared to blue. Platinum-cobalt multilayers, prepared by alternatively sputtering ultrathin films of platinum (1.5 nm) and cobalt (0.3 nm), and thin films of Co-Pt alloys (25-nm thick, 45-90% Pt), made by electron-beam evaporation, show superior behavior (Gurney, 1993; Weller et al., 1992). The binary cobalt-platinum alloy films have a large perpendicular magnetic anisotropy, and show enhanced Kerr rotation and Kerr ellipticity compared to those of Tb-Fe-Co and even to those of platinum-cobalt multilayers. They can be read at the wavelength of 450 nm, have 100% perpendicular remanence, and coercivities of the order of 200 kA/m. Cobalt-platinum layers for magneto-optical storage can, contrary to those made from rareearth alloys, be stored in the open air. They are expected to be the most promising magnetic layer materials for the next generation of storage devices.

Heat Mirrors

Window-glass panes coated with thin transparent silver layers are an effective means against energy losses by radiation in buildings. Caused by the high transparency of the glass in the visible and in the near infrared wavelength range, glass windows, glass doors, and glass fronts are the most important heat flow transition regions for incoming and outgoing radiation. Silver reflects light waves extending from the visible range into the infrared. Thin layers of silver at thicknesses below « 0.012 Jim become transparent to optical light waves in the visible range (380-780 nm), but retain their high degree of reflection in the long wave region. Thin, optically transparent silver coats deposited on architectural window glasses therefore reduce energy losses caused by outgoing heat radiation. Additional thin, transparent, and neutral color oxide coatings serve as antireflection layers and improve the adhesion, corrosion resistance, and mechanical stability. The optical quality is influenced by the structure of the coatings. The silver layers must be coherent and must not show unconnected silver islands or hillocks. Both lower the reflectance and cause losses in the transmittance by scattering the light (Dachselt e t a l , 1982; Glaser, 1989). Gold coatings are particularly used on architectural glass for sun protection. Chemical Technology

Corrosion resistance, mechanical stability, and good thermal conductivity determine the use of noble metals in chemical technology as constituents in materials, for example, as parts or components of equipment, apparatus, crucibles, or as protective surface coatings. Silver has advantageous mechanical and working properties. It finds extensive use

9.5 Applications

in alkaline fusion processes in laboratory equipment. Silver alloys are important constituents for vessels, containers, and similar equipment which has to withstand corrosion by inorganic and organic acids and at the same time requires good thermal conductivity. These parts or components can be made from bulk material or applied as surface coatings, plated from sheet material or by galvanic processes in thicknesses from « 50 jim to 1 mm. Alloying with 0.15 wt.% nickel prevents recrystallization at elevated temperatures. Silverplated parts are used in vessels, autoclaves, pipework, heat exchangers, and fittings. Solid, bulk silver components are used in strongly alkaline and hydrochloric solutions and for fruit juices. Filter elements of sintered spheroidal silver powder are used in the food industry in fruit presses. Silver and gold rings or foils are used for sealing oxygen compressors and for sealing high vacuum equipment at pressures below 1 0 ' 6 Torr (133 x 10" 6 N m~ 2 ). Platinum crucibles, dishes, boats, and electrodes for electrogravimetry are the basic tools for laboratory technology. If dimensional stability is required, components are made from the alloys Pt-Ir3, PtAu5, and Pt-Ir0.3. In the production of highly corrosive materials platinum coatings are also applied, produced either by roll plating, inlay pressing, or electrolytic coating. Porous platinum, palladium, or silver filters are produced for corrosive media by sintering spherical powders. Burst membranes against over-pressure in corrosive atmospheres and for pressure ranges from 1-^400 atm (40 MN m~ 2 ) are made of silver or gold (thermally resistant up to ^1200°C), or platinum or palladium (thermally resistant up to 1400 °C). Diffusion membranes made from palladium for the separation of nitrogen and hydrogen

567

are capable of producing high purity hydrogen ( > 99.999 %). This is of special importance for semipermeable electrodes in fuel cells. Hydrogen diffusion at temperatures below 1000 °C is facilitated by coating the surface with palladium-black emulsions. Resistance heaters for high-temperature furnaces (up to 1500 °C) consist of platinum or platinum-rhodium alloys. Pt-Ni, Pt-Rh, and Pt-Ru alloys are used in spark plug electrodes. Because of the need for corrosion and erosion resistance, platinum metals are preferred for crucibles in the glass industry. High-purity optical glass is melted in pure platinum crucibles. Platinum components (skimmers, pouring funnels, protection tubes for thermocouples, level-measuring devices) ensure problemfree operation in the automatic production of bottles over long periods. Dispersionhardened platinum is also used for equipment parts in glass-working technology. In the manufacture of glass fiber and glass wool, the spinnerets are made of Pt-Rh alloys. Rock wool and slag wool are made by a centrifugal spinning process with platinum centrifuges. Spinnerets for textile fibers (rayon, acrylic) are made from platinum-gold and platinum-rhodium alloys (Fig. 9-42). Platinum-rhodium tubes are used to protect high-temperature thermocouples used in

Figure 9-42. Pt-Rh spinnerets for rayon fibers (Degussa).

568

9 Noble Metals and Their Alloys

glass-melting production. Quartz glass fibers are drawn through iridium spinnerets. Platinum-plated wires of iron, nickel, or molybdenum, or gold-plated iron-nickel wires serve as vacuum-tight leads for glass tubes. Platinum and iridium are indispensible crubile materials for the production of high-purity single crystals for optical and laser technology (Table 955)(Benner etal., 1991, 7).

Table 9-55. NM-containing equipment, components, and crucibles: (a) crucibles for single-crystal manufacture (Benner, 1991, 3), (b) general equipment (Diibler, 1995). (a)

9.5.2.3 Combined Physical and Chemical Interactions

Cz b Gd 3 Ga 5 0 5 ("GGG") (Gd-Ga-garnet)

Catalysis

General Remarks Catalysis is one of the most important technical application fields for noble metals. With the exception of gold, noble metals act as highly efficient and selective catalysts in numerous oxidizing and hydrogenizing reactions of inorganic and organic chemical production processes. Catalysts accelerate chemical reactions by lowering the activation energy. Lowering the activation energy also means that at any predetermined reaction rate, the process can be run at a lower temperature. Catalysts participate in the reaction run, but, after the reaction is terminated, are returned to their original condition. The position of the reaction equilibrium is not affected (Ostwald, 1894). Heterogeneous catalysis is performed with solid catalysts, which are provided either as bulk material (e.g., wire and wire gauzes) or as a surface coat on a suitable carrier ("supported catalyst"). The reactions take place on the catalyst's surface. In homogeneous catalysis, the reactants and the catalyst, which is usually a complex metal-organic compound, are both present in liquid, preferably organic, solution.

Single-crystal Manufacmaterial turing method Ferrite A12O3

Ba Cz b

Use

Noble metal

magnetic heads Pt,Pt-Rh Ir substrates, hybrid IC bearings, jewels Ir

bubble memory substrate jewels

LiTaO3

Cz b

surface elastic Ir, Pt-Rh wave elements

LiNbO 3

Cz b

surface elastic wave elements

a

Bridgman method, (b)

b

Pt

Czochralski method,

Equipment

Noble metal

crucibles

Pt, PtAu5, PtRhlO

bowls

PtAu5

Burn to ash

crucibles, dishes

Pt, Ptlr3 AuPtlO, AuPd20

Evaporation, drying

bowls

Pt, Ptlr3

electrodes

Pt

Use Melting or solution of minerals and ferrites X-ray analysis

Electroanalysis Gravimetry

net electrodes

PtMO

Gas chromatography

tubules, cannulas

Ptlr25

Viscosimetry

crucibles, stirrer

PtRh(10-20) PtAu5

porous discs

Ag, Pt

Filtration

9.5 Applications

Heterogeneous Catalysis In heterogeneous catalysis, action takes place on the catalyst's surface on "active centers". The main reaction steps are: • adsorption of at least one reactant, • reaction of adsorbed species to an intermediate product, • reaction of the intermediate product with another reactant to the final product, and • desorption. Temporary bonding of the reactant molecules on the surface of the solid catalyst alters the forces between the adsorbed atoms and changes the energetic requirements for the subsequent reaction. Bonding can be caused by the transfer of electrons from the catalyst's surface to the adsorbed molecules, weakening their molecular bond strength and facilitating their dissociation. The specific catalytic action of the catalyst depends on its electronic structure, atomic size, and surface structure. A necessary condition for the capability to adsorb the reactants and to desorb the products is given as a moderate bond strength between the adsorbed species and the catalyst's surface (Tamara and Naito, 1991; Yates 1992). The molecules are adsorbed in either molecular or in dissociated form. Because of the surface-induced reaction, high specific surfaces (usually between 100 and 1000 m2/g) are favorable for higher yields and reaction rates. Highly dispersed noble metal coatings are precipitated on catalyst supports by thermal decomposition or chemical reduction of noble metal salts ("platinum-black", "palladium-black"). Catalysts can become deactivated ("poisoned") if the active surface areas are occupied by atoms or molecules with form strong bonds. The most dangerous contact poisons for PGM catalysts are

569

sulfur, phosphorus, carbon monoxide, hydrogen cyanide, lead, and mercury. The heat of adsorption of gaseous molecules of O2, H 2 , N 2 , CO, and CO2 decreases in each period of the periodic table of the elements with increasing atomic number, and reaches medium values for the group 8 elements. Based on the bond strenghts of the diatomic molecules (CO>N 2 >NO), dissociative adsorption takes place at elements with the corresponding dissociative capacity. The heat of adsorption on different crystal planes of f.c.c. crystals increases with decreasing density of surface atoms in the order {lll}>{100}>{110}. Correlations also exist between the lattice dimensions of catalysts and the dimensions of adsorbed molecules. The catalytic activities for the hydrogenation of C = C double bonds decrease in the order R h > R u > P d > P t . The key step is the dissociation of hydrogen molecules to atoms on the surface. The lattice constant of rhodium is closest to the distance of the hydrogen atoms between the methyl-groups of the ethane molecule (Schuit and Van Reijen, 1958; Boudart, 1977). The order ranges of effectiveness of the PGMs for structure-sensitive catalytic activities differ for different reaction types, for example, 1) for the oxidation of paraffins (methane, etc.): Pd>Pt>Co 3 O 4 >PdO; 2) for the "methanation" of CO with H 2 : CH 4 synthesis activity: Ru>Rh>Pd>Pt>Ir, CO dissociation ability: Ru>Rh>Pt>Pd; 3) for the synthesis of paraffins [FischerTropsch (Vannice, 1975, 1977)] according to nCO + 2nH2

+ nU2O

(9-22)

570

9 Noble Metals and Their Alloys

with Ru>[>Fe>Ni>Co]>Rh>Pd > Pt>Ir; 4) for the reduction of NO with CO: Pt > Pd > Rh > Ru; and 5) for the reduction of NO with H 2 : Ru>Rh>Pt>Pd (Obuchietal., 1982, 1983). The selectivity of catalytic reactions and the long-time stability of the catalysts can be improved by alloying. Platinum/rhenium catalysts have improved yield and improved temperature stability compared to pure platinum catalysts in reforming reactions (Pollitzer, 1972; Burch, 1978). Iridium also acts as a promoter in reforming processes. Further examples of the modification of the catalytic selectivity are given in Table 9-56.

In the catalytic reaction for the formation of hydroxylamines, which are the starting material for nylon, 2 NO + 3 H, >2NH2OH

(9-23)

with a platinum/carbon catalyst, formation of the unwanted by-product of ammonium sulfate is suppressed by "partial poisoning" of the catalyst with mercury or silver salts. Additions of halogens prevent unwanted sintering of platinum particles, minimize adhesion of carboneous substances to the catalyst's surface, and participate in the redispersion of the platinum during regeneration (Benner 1991, 3). Metallic silver is an excellent catalyst for oxidation processes of organic compounds. It is used in the form of wire gauzes or precipitated on A12O3 or silicate carriers. Technically important processes

Table 9-56. Modification of catalytic activities by alloying (Renner, 1992). Additional metal

Reaction

Effect of additive

Pt

5-20% Rh

ammonia oxidation

increased NO yield lower Pt losses

Pt

Ge, Sn, In, Ga

dehydration and hydrocracking of alkanes

longer catalyst life due to fewer deposits

Pt

Sn + Re

dehydrocyclization and aromatization of alkanes

increased catalyst activity and stability

Pt

Pb, Cu

dehydrocyclization and aromatization of alkanes

effective aromatization

Au

oxidative dehydrogenation of alkanes, ra-butene to butadiene, methanol to formaldehyde

better selectivity

Cu(Ag)

reforming

less hydrogenation relative to isomerization

Ir

Au(Ag, Cu)

hydroforming of paraffins and cycloalkanes

high yield of aromatics above 500 °C

Pd

Sn, Za, Pb

selective hydrogenation of alkynes to alkenes

Pd

Ni, Rh, Ag

alkane dehydrogenation and dehydrocyclization

Base metal

Pt, Pd, Ir

Ru, Os

hinders coking

9.5 Applications

571

are the production of ethylene oxide (Lefort, 1931) H

H C=C

H

H

o o

H

x -C c

1 \

\l

H/ ^\

H H

+ O"

|

urn

O (9-24) and the catalytic dehydrogenation of methanol to formaldehyde. This catalytic reaction is supposed to proceed as a surface reaction between the temporarily adsorbed ethylene and oxygen molecules (Campbell, 1984). The platinum group metals show pronounced catalytic action in hydrogerfation and oxidation reactions (e.g., reforming processes in petrochemistry, the oxidation of NH 3 to NO, hydrogenation of hydrocarbon compounds, etc.). PGM catalysts are also used in wire gauze form (Fig. 9-43)v or as metal-black surface layers on ceramic supports (Fig. 9-44). The noble metals are precipitated from their chloride or nitrate solutions onto the supports of e.g., carbon, A12O3, silica gel, aluminum silicate (zeolite), or CaCO 3 . The support material can be in the form of powder or granules; the concentration of noble metals varies between 0.5 and « 5 % . The largest demand for platinum catalysts in the chemical industry comes from petroleum refineries, especially for the reforming of high boiling fractions from the distillation of crude oil at atmospheric pressure. The main products are gasolines with good antiknock properties. The catalysts used are Pt-Al 2 O 3 , Pt-Ir, and Pt-Re. Further uses are in hydrorefining and hydrocracking processes. The pharmaceutical industry uses combinations of active carbon supports with palladium, rhodium, and platinum, preferably for hydrogenation processes. A typical example of solid metal catalysts are fine-mesh gauzes of 60-|im wire of

Figure 9-43. Pt-Rh catalyst gauze (Degussa).

Figure 9-44. Platinum-black catalyst coating (Degussa).

Pt-RhlO, used for the large-scale oxidation of ammonia to nitric acid (Ostwald and Brauer, 1931; Sperner and Hohmann, 1976), and in the production of hydrogen cyanide (Andrussow, 1935/55). The modern (BMA [German abbreviation for Blausaure aus Methan und Ammoniak]) process (Endter, 1958) uses A12O3 tubes which

572

9 Noble Metals and Their Alloys

Table 9-57. Catalytic processes in chemistry (DEGUSSA, 1967, Prescher, 1995). Noble metal catalyst

Reaction

Powder a) Oxidation and other reactions on noble metal catalysts Methanol -• formaldehyde Ag Ethylene -» ethylene oxide Ethylene -• acetaldehyde Amines -» nitriles Complete oxidation of organic compounds

RuO 4 in CC14 -

Rh, Os

Ru, Pd

NH 3

PtRhlO

b) Reduction and isomerization on noble metal catalysts H 2 addition to acetylene and homologs

ethylene and homologs aromatic hydrocarbons aromatic compounds, hydrogenation of the aromatic ring heterocycles, hydrogenation of the aromatic ring Reforming of crude oil dehydration (cracking) to aromatics isomerization/cyclization cleavage to alkanes

Ag/Al2O3 PdCl2 in solution Pd/Al 2 O 3

Inorganic catalysis H2O2 - H2O + J O 2 NO -• hydroxylamine

NO

Support

Ag

Disulfides to sulfoxides Alcohols to aldehydes, ketones Olefins, oxidative cleavage Synthesis or decomposition of organic compounds CO + H 2 to methane to paraffin hydrocarbons Addition of CO and H 2 to olefins to give alcohols olefins to give ketones olefins to give esters Addition of silane to olefins Cleavage of formic acid Cleavage of formaldehyde

Compound

RuO 2

Ru/Al 2 O 3 Ru + K 2 CO 3

RhO 3 RhCl3 Rh oxide Ru oxide OsO4

_ Pt/carbon A12O3 Pt/asbestos Pd/BaSO4

Pt PtAu5 ^/carbon Ptlrl Rh

Rh, Pd, Pt

-

-

-

Rh, Pd ") ( Ag/Pd70/30J/ 2< Pd and Pt/BaSO4 Ru/C Rh/Al 2 O 3

RuO 2

Rh/Al 2 O 3 Rh/Al 2 O 3 Pd | Pt >/Al2O3 Ptlr J Ru/SiO2

9.5 Applications

573

Table 9-57. (continued) Reaction

Noble metal catalyst Powder

Compound

Support

of carboxylic acids to alcohols

RuO 2

Ru^j Pd f /carbon Pt 3 Ru/carbon

nitrogen compounds to amines

PtO 2

of disulfides to mercaptans of mercaptans to alcohols + H 2 S

RuO 2 RuO.

Reduction of ketones to alcohols

are coated on the inside with a platinum layer. Coatings of palladium-black are used as catalysts for the production of hydrogen peroxide from air and hydrogen. Table 9-57 gives a survey of the most important chemical processes which are performed with noble metal catalysts. Automobile exhaust gas catalysts take the biggest portion of the total demand of platinum (44%) and rhodium (83%). They reduce and remove the gaseous and particulate pollution constituents from exhaust gas streams by catalytically activated reduction and oxidation reactions. Different types of catalysis have been developed for gas and diesel engines. In gas engines, CO and hydrocarbons (CHX) have to be oxidized to carbon dioxide and water, and nitric oxides (NOX) have to be reduced to N 2 . The exhaust gases from diesel engines have to be cleaned of carbon and adsorbed liquid particulates by catalytic oxidation (Taylor, 1993). The three-way catalyst for gas engines contains as the active part a highly dispersed metallic coat of a mixture of platinum, 20 wt.% rhodium and palladium in changing proportions. The metal coat is precipitated on a honeycomb-shaped ceramic ("monolith") support made from

Pd f /carbon Pt

cordierite (Fig. 9-45). The monolith is coated with a ceria-doped "washcoat" of A12O3, which forms a porous structure with a highly specific surface area on firing. The oxidation of CH^ is activated by platinum and palladium, and the reduction of NOX to nitrogen by rhodium. The reaction takes place in consecutive steps. When CO and NO meet the catalyst surface, both are adsorbed at different places. The N - 0 bond is weakened by adsorption on the rhodium surface, thus enabling dissociation. The nitrogen atom stays in its place while the dissociated oxygen atom combines with the adsorbed CO molecule to form CO 2, and desorbs. In the second step, another adsorbed nitrogen atom combines with the prior one to form an N 2 molecule, and desorbs (Friend, 1993; Cho, 1994; Madix, 1986). The reaction has its best yield at an air to fuel ratio of 14.7:1. The ratio is measured and maintained by means of two platinum sensors. One is installed in the inlet manifold to monitor the proportion of air in the air-fuel mixture as it enters the combustion chamber, the second forms the tip of the so-called lambda probe, inserted in the exhaust stream ahead of the catalyst to measure the amount of oxygen in the

574

9 Noble Metals and Their Alloys

(a) r

Control function

'

•

i

i

J X-probe r. -—:—+ Air Mixer Petrol

1 1 I Petrol-motor

(1

i

N

Fhreeway :atalytic zonverter

1 J Nr I/

(b)

(c)

Figure 9-45. Three-way catalytic converter: (a) schematic view (Prescher, 1995), (b) cut-away view (Johnson Matthey, 1992), and (c) Pt/Rh catalyst surface layer precipitated on the washcoat (Prescher, 1995).

575

9.5 Applications

exhaust. Each catalytic converter contains « 2 g PGM mixture of either platinum/ rhodium or platinum/palladium/rhodium in varying proportions. Three-way catalysts of the same type, together with lambda probes are also used in thermal power stations. Carbon particulate emissions of diesel engines are removed by passing the exhaust gas stream along heated platinum- and palladiumcoated ceramic filters. The carbon particles precipitate on the surface and are burned.

and the formation of elastomers and products of silicone chemistry. The synthesis of acetic acid from methanol (Monsanto process) yO

CH3OH + CO -> H 3 C - c'

(9-27) X

OH

is carried out in the presence of soluble rhodium catalysts. The Wacker process for the synthesis of acetaldehyde HN

(9-28) N

Homogeneous Catalysis

H'

In homogeneous catalysis the catalysts and the reactants are present in the same (liquid) phase. Compared to heterogeneous catalysis, homogeneous catalysts show much higher selectivity, and enable milder reaction conditions and higher yields. Homogeneous catalysis is used in organic chemistry in many technical processes for the production of detergents, and softening agents for plastics and pharmaceuticals. Complex rhodium compounds of the tpye [RhCl (PPh 3 ) 3 ] (here PPh 3 is triphenylphosphate), "Wilkinson's catalyst", (Osborne et al., 1965; Coffey, 1966) are especially suitable for the homogeneous catalysis of hydrogenation and hydroformylation reactions. Typical reaction types include hydrogenation

uses PdCl 2 or H 2 PdCl 4 (Smith, 1975). Table 9-58 shows some examples of precious metal compound catalysts used in homogeneous catalytic reactions. Homogeneous catalysis is especially important for the synthesis of complex enantiomer pure organic compounds. Chiral noble metal complex compounds effect stereospecific chemical synthesis. The most important catalysts are based on rhodium. Special interest is given to the synthesis of enantiomer pure chemicals which have major importance as Pharmaceuticals. Platinum metals are effective catalysts in the photodissociation of water by lightdriven electron transfer. Three different re-

r

H-C-C-H

(9-25)

the large-scale hydroformylation ("oxo"process) for the production of aldehydes from olefins, carbon monoxide, and hydrogen CO/H 2

I H H- C- C S O

O

Table 9-58. PGM compounds for homogeneous catalysis (Prescher, 1995). [RhCl(Pph3)3]a [RhH(CO)(Pph3)]a [Rh(Nbd)(Dppb)+ BF 4 ] c ' d [RuCl2(Pph3)3]a [IrCl(CO)(Pph3)3]a [Ir(Cod)(py)P(C 6 H 11 ) 3 ] + PF- b ' e a

(9-26)

H

ph=phenyl; b py=pyridine; c Nbd=norbornadiene; Dppb = 1,4-diphenylphosphinobutane; e Cod = 1,5cyclooctadiene. d

576

9 Noble Metals and Their Alloys

action types are under development (Mills, 1991): 1. The homogeneous reaction type uses complex metal-organic compounds of ruthenium, rhodium, or zinc as photosensitizers to promote electron transfer between suitable electron mediators such as electron acceptors and electron donors. Compounds in this category are of the polypyridyl Ru11 complex type: [Ru(bipy)3]2 + , which act both as electron acceptors and electron donors. If electrons are excited by visible light, the compound has sufficient reduction potential to decompose water in the presence of suitable catalysts (Watkins, 1978). Hydrogenevolving systems use colloidal platinum suspensions as catalysts. Increased yields are obtained by alloying with 20-30 at.% gold (Harriman 1991). The photogeneration of hydrogen has also been performed with reactants which are loaded in special permeable membranes (solid polymer electrolyte membrane = SPE system). The catalyst components are TiO 2 particles coated with platinum, the electroactive compounds are polypyrrole and chlorophyll (Khare et al., 1994). 2. The photo-induced charge-separation in inorganic semiconductors is based upon irradiation of inorganic semiconductors with light of higher energy than the band gap of their electrons. Irradiation results in the formation of an electron/hole pair, in which the electron resides in the conduction band and the positive hole remains in the valence band. These electrons and holes can migrate separately to the semiconductor surface forming hydrogen and oxygen under the influence of PGM group catalysts. Optimum, highly effective semiconductor materials have band gaps of ^2.2 eV. Promising results have been achieved with CdS (band gap 2.4 eV),

which has been stabilized by loading with 1 wt.% RuO 2 (Harriman, 1983), and with lower band gap semiconductor crystals (silicon) coated with thin platinum layers (Maier and Bilger, 1994).

Electrochemistry Fuel Cells

Fuel cells convert chemical energy directly into electrical energy by catalytic oxidation of hydrogen. They are battery-type devices, consisting of two complex catalyst-bearing electrodes, which are separated by an ion-conducting electrolyte (Fig. 9-46). Hydrogen is oxidized at the anode and oxygen is reduced at the cathode, generating an electric current together with some reaction heat. Fuel cell technology can contribute to a considerable extent to the conservation of natural resources, operating at much higher electrical efficiency during the energy conversion process (up to 60%) than conventional power plants. Several different fuel cell types have been developed, operating with different electrolytes and at different temperatures. Noble metal catalysts are used for devices operating at temperatures below 200 °C (Table 9-59). At present, two different types: the phosphoric acid fuel cell (PAFC) and the proton exchange membrane fuel cell (PEM) are the most advanced. PAFC devices are preferably designed for high power generator units ranging between 50 kW and 200 kW, mainly for the on site generation of electric power and heating. The electrolyte is phosphoric acid. The electrodes are composed of supporting coal fiber on the gas side and a multicomponent porous structured layer made from a mixture of platinum/carbon catalyst and PTFE [poly(tetrafluoroethylene)]

9.5 Applications

577

Table 9-59. Fuel cell types (Johnson Matthey, 1993), Type

Air

Fuel—— H

Working temperature (°C)

KOH

50-90

polymer, e.g., Nation (DuPont)

50-125

H 3 PO 4

190-210

eutectic melt of Li 2 CO 3 /K 2 CO 3 ZrO 2 , doped with yttrium

630-650 900-1000

rent density. They operate at temperature ranges of 50-125 °C. A thin, solid membrane of an ion-conducting polymer serves as the separating, proton-exchanging electrolyte. The electrodes are directly coated on both sides, also using PTFE-containing porous layers with platinum/carbon on the cathode and platinum (plus ruthenium and palladium)/carbon on the anode side. Bifunctional cells, which generate electricity either by hydrogen oxidation or hydrogen production by electrolysis may become of interest for temporary storage of excess electrical energy from photovoltaic plants.

Hydrogen H

(b)

Alkaline fuel cell (AFC) Proton exchange membrane or polymer electrolyte membrane (PEM) Phosphoric acid fuel cell (PAFC) Molten carbonate fuel cell (MCFC) Solid oxide fuel cell (SOFC)

Electrolyte

Polymeric electrolyte

Product water H 2 0 + 02

Figure 9-46. (a) Principle of a fuel cell (schematic diagram), (b) Principle of a polymer electrolyte membrane fuel cell (StraBer, 1990).

on the electrolyte side. Cathode catalysts are Pt/Co/Cr, Pt/Co/Fe, or Pt/Ni alloys. PEM type fuel cells are especially suited for smaller units, e.g., for electrically powered vehicles, because of their higher cur-

Electrodes Large quantities of noble metals are used as electrode coating materials. Coatings of elementary PGMs, PGM mixtures, or PGM oxides exist. The base material is in most cases titanium in the form of sheet, tubes, or expanded metal. Costings are applied either by electroplating or by applying mixtures of noble metal chlorides with alcohols and terpene oil on the surface, drying and firing at 550 °C to decompose

578

9 Noble Metals and Their Alloys

and form the active deposit. The most often applied type is the platinum-plated titanium anode, used in many different applications like seawater electrolysis, high speed tin-plating of steel (e.g., food canning), and general purpose electrodes in cases where high purity of the galvanic bath and exact and unchanging distances between the anode and the cathode are required. RuO 2 and platinum or Pt-IrO 2 coated electrodes have a lower overpotential against chlorine and therefore find use in alkali-chlorine electrolysis, saving energy and giving controlled composition of the chlorine gas (low O 2 concentration). IrO 2 electrodes are also used in desulfurizing plants for flue gas. In electronic production processes with highly sophisticated plating equipment, platinum-plated titanium electrodes are used for plating gold with noncyanidic bath compositions, especially for plating lead frames, semiconductor packages, contact pads, or for continuous selective or spot-plating. Batteries Silver is the only noble metal used in batteries, applied as silver oxide for cathodes in combination with either zinc, cadmium, or magnesium as the anode. The AgO/Zn button battery type is of great practical importance as a current source for cameras, watches, hearing aids, and similar small devices. The worldwide yearly demand exceeds more than one billion pieces. Ag-0 batteries have high energy densities (80-250 W/dm 2 ), high discharge current, and high long-life stability. The overall reaction

II)

(9-29) OH) 2 (9-30)

develops a cell voltage of 1.55 V. For rechargeable units, cadmium is the preferred anode material. AgO-Cd cells have a lower but constant cell voltage. In recharging cycles they reach lifetimes exceeding 1000 cycles (Fleischer and Lander, 1971; Lingen, 1983). AgO-Mg is used in seawater batteries, in which the salt content of the seawater represents the electrolyte. These units operate automatically and they are important in sea rescue operations. Photosensitive Materials Noble metals play an important role as constituents of photosensitive compounds. Silver is used in large quantities (at present «40 % of the total worldwide yearly demand) as halide compounds. In all cases, high purity silver nitrate is the starting material, while gold, platinum, iridium, rhodium, and ruthenium salts serve in small quantities of parts per million (ppm) as sensitizers and additives to the silver halide matrix. Photosensitivity and image formation are principally based on changes in the oxidation states of the noble metal cations. Light absorption causes excitation of the electrons, which form stable silver clusters with the silver ions (Agn, where n > 4). These act as catalysts for the image-forming reaction during the development. AgBr has Frenkel type lattice defects with interstitial silver ions and vacant silver ion lattice sites. This is favorable for the photolysis of silver bromide forming neutral silver atoms and the silver clusters. Light signals are amplified by this by a factor of ^10 8 , one cluster of four silver atoms reducing ^ 1 0 9 silver ions by catalytic action to stable metallic silver particles. Practical silver film material consists mainly of AgBr, mixed for high speed ma-

9.5 Applications

579

Table 9-60. Photosensitizing NM compounds (Tanaka, 1991, 2). Role

Element

State

Function

Photosensor (and imageforming material)

Ag

AgX microcrystals (X: halogen)

catalyze Ag4 cluster formation by photolysis, Ag + source for black Ag image formation, oxidizing agent for color image formation

Sensitizer

Ag

Ag2 clusters Ag2S clusters

traps for positive holes concentration centers for photoelectrons

Au

AgAuS clusters

Ir

IrXj- (Ir 3+ in Ag + lattice sites)

reinforcement of concentration centers for photoelectrons stabilizer and activator of Ag3 clusters deep temporary traps for photoelectrons

Contrast enhancer

Rh

RhX^~ (Rh 3+ in Ag + lattice sites)

effective recombination centers for photoelectrons and positive holes

Development nuclei

Ag, Au, Pt, Pd, etc.

clusters Ag4, Ag2Au, etc.

catalyzes metal formation by reduction of metal cations

Development accelerator

Ru

Toning agent

Au, Pt, etc.

Ru(NH 3 )| + , etc. chloro-complexes

terials with a few mole percent Agl or, for low speed, high image quality, with a few tenths of mole percent of AgCl. Crystals are prepared by precipitation out of aqueous solutions in the presence of gelatine as a protective colloid. Under appropriate precipitation conditions, platelike-shaped crystals are formed. They have to be sensitized by reducing treatment or the action of sulfide ions to avoid dispersion of photolytically formed silver clusters. Gold additions enforce sensitivity, the gold ions being incorporated into the silver sulfide clusters as Ag 3 AuS 2 and AgAuS. Iridium, incorpoated as Ir m in quantities of 0.01-100 jimol is another effective sensitizer (up to 100 times) for AgBr, especially advantageous for laser light induced photographic processes. Rhodium, incorporated as the octahedral complex in RhCL 3 ~, acts as a contrast enhancer.

electron-transfer mediators conversion of black Ag image to Au, Pt, etc.

Table 9-60 gives a survey of the main types of noble metal compounds and their action in photosensitizing materials (Tanaka, 1991). In the development of black-and-white film and paper, the silver halide grains, which have been exposed to photons, are preferentially reduced by chemicals to elemental silver. Unexposed grains are dissolved by sodium thiosulfate, leaving the black silver deposits as stable picture. In color photography, the three primary colors are each produced by a separate photo-layer, sensitized for each color by special dye components. Light exposure generates primarily a silver image, but corresponding to the sensitized color of each layer. These silver deposites activate formation of the colored picture. The deposited silver is "bleached" by oxidizing agents and dissolved together with the unexposed

580

9 Noble Metals and Their Alloys

halide components in the thiosulfate dissolution treatment. The ready colored picture does not contain silver. Silver is recovered from the used solutions by special, complex chemical processing.

9.6 Recycling The recovery of noble metals from industrial residues, scrap, used and reproducible material ("secondary materials") is of major economic importance regarding both their high commercial value and the limited reserves. Recycling includes controlled collection, separation of material according to composition, purity, noble metal concentration, and pure metal compounds (oxides, dross, and slag). Recycling processes are essentially the same as those for the production of primary materials, including chemical, pyrometallurgical, hydrometallurgical, and electrochemical processes. Figure 9-47 gives a schematic overview of the main processing routes for recycling precious metals. 9.6.1 Direct Melting Properly selected, pure metallic scrap can often be directly remelted to metal alloy anodes for further processing. This is applied for high grade Au-Ag and PGM alloys which can be fed into the converter furnace (Fig. 9-48). Gold can be extracted from alloys containing more than 30% gold by passing chlorine gas over the surface of the liquid alloy (Miller process). 9.6.2 Dissolution High grade PGM-containing metallic waste from used equipment and semifinished products like crucibles, dishes, thermocouple elements, gauze catalysts, and fiber spinneret nozzles are dissolved with-

out prior treatment in HC1/C12 or aqua regia for further separation by hydrometallurgical processes. Direct dissolution can also be applied to gold-rich alloys. Silver is precipitated as AgCl, gold by selective reduction with SO2, hydrazine, or, if PGMs are present, with oxalic acid. Gold and silver coatings are removed from substrate material by alkaline cyanide solutions. From these, gold is precipitated with zinc powder; silver is recovered by electrolysis. Base metals can also be removed from silver alloys by acid leaching or scrap-metal electrolysis. 9.6.3 Pyrometallurgical Processes and Equipment Pyrometallurgical processes are applied to low grade and compound scrap material for extraction, separation, and concentration to high noble metal content. Extraction is performed by melting with "collector metals", dissolving the major portion of noble metal. The slags, oxides, and base-metal alloys, containing lower amounts of noble metals, are separately treated or fed back to the preceding operation. The collector metal for silver and gold is lead; the collector metal for essentially PGM-containing materials is copper. Figure 9-48 shows schematically the recycling process for silver scrap. In the 'blast furnace' running at temperatures between 1100 and 1300 °C, PbO is reduced by CO and the liquid lead extracts the silver from the scrap together with minor amounts of accompanying metals (gold, PGMs, base metals) up to a concentration of ^ 2 0 % . The silver-containing lead bullion is passed to the converter furnace for oxidizing treatment to remove base metals by forming liquid PbO and upgrading the remaining silver alloy to ^ 9 8 % silver. Working temperatures rise in relation to

581

9.6 Recycling

Metallic* nonrnetallic

Metallic starting scrap material Homogenizing sampling Concentration separation

High gold content ( + A g < p G M ( basemetal)

High Ag content (+Au, PGM, basemetal)

Noble metal mix (+dross, basemetal)

PGM

Melt

Melt

Preparation

Melt

Ag, PGM |

Converter

basemetal

furnace

PGM+ceramic, basemetal, support catalyst autocatalyst Preparation

Dross basemetal

PbO Miller process

Metallic* nonrnetallic

Metallic

Retining

Precipitation — + - P(j

Figure 9-47. Recycling of NM scrap (schematic diagram).

Precipitation — » -

Blast furnace

Converter

Collection of noble metals by extraction with molten lead

Removal of base metals (Pb, Cu, Bi, Sn, etc.)

Reduction process: PbO + CO — • Pb + C02

Oxidation process:

Dross (containing noble metals)

Additives Coke

Off-gas

Litharge

Reaction zone

Air Slag Copper matte Lead bullion

Scrap metal (containing noble metals)

PbO Auriferous silver (98%Ag, 2%Au) Electrorefining

Lead bullion (Pb, 20%Ag) Silver 99.99%

Figure 9-48. Converter furnace (Renner, 1993, 3).

582

9 Noble Metals and Their Alloys

the increasing silver content of the melt from 900-1100 °C. The liquid PbO (litharge) with 5 % silver is continuously fed back into the blast furnace, where it is again reduced to metallic lead. The liquid Ag-Au alloy is cast to anodes for further separation and refining. Copper melts with high-grade PGM material are only partially oxidized ( ^ 2 0 % of the melt) to remove scrap and base metals, because at higher PGM concentration in the melt, the Cu 2 O would take off considerable quantities of PGMs and reduce the process yield. High grade, pure metallic scrap can be directly melted in the converter for further treatment. A special furnace type ("submerged arc furnace") is heated by direct current flow through the bottom, using the metal/scrap load as resistor material, connected with a counter electrode. The method is especially suitable for the recycling of PGM automotive exhaust catalysts, which contain high portions of A12O3, by extraction with copper. After formation of the melt, additional scrap material can be fed through the tube-shaped counter electrode. PGMs are concentrated up to 20%, the melt separated and cast to anodes for separation and purification by hydrometallurgical processes. Anion exchange resins are used to extract gold, platinum, and palladium from highly diluted solutions. Solvent extraction offers the possibility of separation with a quantitative yield and a highly selective action to be applied for gold, platinum, palladium, and iridium. It works with special liquid organic compounds which form complex or ion-pair compounds with the noble metals in aqueous solution. These dissolve predominantly in the organic solvent and can be separated if the organic and aqueous phases are not miscible. Purification and reduction to

metals is done in the same processes used for recovery from aqueous solutions. The recovery of silver from photographic material (film, fixing liquids) includes cutting, burning, or leaching of the organics, and further refining by alternating chemical solution and precipitation and electrolytic processes. Special care has to be taken to suppress the silver concentration of wastewater below 0.04 ppm, because of harmful effects on biological systems. The spent fixing baths that contain approximately 5 g/1 silver, are treated by reduction electrolysis. Exposed films, for example, expired X-ray films are incinerated or treated by pyrolysis. Alternatively, the silver grains in the gelatine layer can be oxidized using iron(m) chloride solution, and the silver halide dissolved with fixing bath solution. The gelatine layer is washed out, leaving elemental silver as well as silver halides which can easily be reduced to metallic silver. Further refining is done in the usual way. The recovery of noble metal plating waste liquidus is done by electrolysis, also by ion exchange processes or by chemical reducing agents.

9.7 References Ahlborn, H. (1965), Z. Metallkd. 56, 205. Ahlborn, H., Wassermann, G. (1963), Z. Metallkd. 54, 1. Ahlborn, H., Wassermann, G. (1964), Z. Metallkd. 55, 167, 685. Andrussow, L. (1935), Angew. Chem. 48, 593. Andrussow, L. (1955), Chem.-Ing.-Tech. 27, 469. Bardeen, J., Cooper, L., Schrieffer, I (1957), Phys. Rev. 108, 1175. Baving, H. X, Emmerich, R., Faulhuber, X, Hartmann, D., Kast, D., Lange, E., Schulten, R. (1991) in: Metall Forschung undEntwicklung, ed. Degussa Ressort Metall, Degussa, Frankfurt/Main, 1: p. 193; 2: p. 198; 3: p. 196. Benner, L. S., Suzuki, T., Meguro, K., Tanaka, S. (Eds.) (1991), Precious Metals Science and Technol-

9.7 References

ogy, Allentown PA, Int. Prec. Met. Inst. (IPMI), 1: p. 636; 2: p. 628; 3: p. 472; 4: p. 725; 5: p. 697; 6: p. 541; 7: p. 336. Bernstein, L., Bartholomew, H. (1965), /. Met. 703. Beuers, X, Krappitz, K. (1995), in: Edelmetalltaschenbuch, 2nd ed. Frankfurt/Main: Degussa. Bohm, H. (1968), Einfuhrung in die Metallkunde: B.I. Hochschultaschenbiicher, Band 196. Mannheim: Bibliogr. Inst. 1: p. 44, 2: p. 103. Both, A., Lollmann, M., Sigel, Th., Zilske, W. (1991) in: Metall Forschung und Entwicklung, ed. Degussa Ressort Metall, Degussa, Frankfurt/Main, p. 59. Boudart, M. (1977), in: Proc. Vlth Int. Congress on Catalysis: Bond, G. C , Wells, P. B., Thompkins, F. C. (Eds.), Vol. 1, p. 1. Brewer, L. (1963), University of California Report UCRL 10701. Brewer, L. (1965), University of California Report UCRL 10701. Brewer, L. (1967), Ada Met. 15, 553. Brodsky, M. B., Marikar, P., Friddle, R. I, Singer, L., Savers, C. H. (1982), Solid State Commun. 42, 675. Brunauer, P., Emmett, H., Teller, E. (1938), J. Am. Chem. Soc. 60, 309. Burch, R. (1978), Plat. Met. Rev. 22(2), 57. Butts, A., Coxe, C. D. (1967), Silver, Van Nostrand, p. 111. Cahn, R. W., Haasen, P. (1983), Physical Metallurgy, 3rd ed. Amsterdam: North-Holland. Cameron, D. S. (1990), Plat. Met. Rev. 34(1), 26. Campbell, C. T. (1984), J. Vac. Sci. Technol. A2,1024. Carcia, P. R, Zeper, W. B. (1990), IEEE Trans. Magn. 26, 1703. Cava, R X, Takagi, H., Batlogg, B., Zandbergen, H. W, Krajewski, J. X, Peck W. R, Jr, van Dover, R. B., Felder, R. X, Siegrist, T., Mizuhaski, K., Lee, X O., Eisaki, H., Carter, S. A., Uchida, S. (1994), Nature 367, 146. Cho, B. K. (1994), J. Catal. 148, 697. Christensen, N. E. (1972), Phys. Status Solidi (b) 54, 551. Christensen, N. E., Saraphin, B. O. (1971), Phys. Rev. B4, 3321. Christie, I. R., Cameron, B. P. (1994), Gold Bull 27(1), 12. Claus, C. E. (1845), Ann. Phys. (Poggendorf) 64, 622. Coffey, R. S. (1966), Br. Patent 121642. Colbus, X, Zimmermann, K. F. (1974), Gold Bull.

7(2), 42. Coombes, X S. (1991/92/93), Platinum. London: Johnson Matthey. Coxe, C. D., McDonald, A. S., Sistare, G. H., Jr. (1979), in: Metals Handbook, 9th ed., Vol.2: Metals Park, OH: ASM, p. 671. Dachselt, W. D., Miinz, W. D., Scherer, M. (1982), Proc. SPIE 324, 37. Darling, A. S. (1967), Plat.Met. Rev. 11, 138. Darling, A. S. (1969), Plat. Met. Rev. 13, 53. Darling, A. S. (1970), Plat. Met. Rev. 14 (1), 15.

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Darling, A. S. (1973), Some Properties and Applications of the Platinum Group Metals: London: Institute of Metals, 1: p. 93; 2: p. 112; 3: p. 117; 4: p. 96; 5: p. 103; 6: p. I l l ; 7: p. 94; 8: p. 101. Degussa (1994), Edelmetallmdrkte. Degussa: Frankfurt/Main. Dermann, K., Groll, W, Kump, U. (1991), in: Metall Forschung und Entwicklung, ed. Degussa Ressort Metall, Degussa, Frankfurt/Main, p. 147 Diehl, W. (1995), in: Edelmetalltaschenbuch, 2nd ed. Frankfurt/Main: Degussa. Diehl, W, Drost, E., Fehrenbach, G.-W., Hohenstatt, M., Huck, R., Jacques, H., Platen, W., Schafer, W, Scheubel, W., Stoll, W. (1991), in: Metall Forschung und Entwicklung, ed. Degussa Ressort Metall, Degussa, Frankfurt/Main, 1: p. 107; 2: p. 114. Doduco Datenbuch (1977), Pforzheim: Doduco, 1: p. 20, 2: p. 50, 3: p. 67, 4: p. 67. Dorbath, B., Frischkorn, H.-X, Jeske, G., Starz, K. A. (1991), in: Metall Forschung und Entwicklung, ed. Degussa Ressort Metall, Degussa, Frankfurt/ Main, 1: p. 45; 2: p. 47, 3: p. 56; 4: p. 57. Drost, E., Blass, A. (1995), in: Edelmetalltaschenbuch, 2nd ed. Frankfurt/Main: Degussa. Drost, E., Hausselt, X (1992), Interdiscip. Sci. Rev. 17, 271. Diibler, K. (1995), in: Edelmetalltaschenbuch, 2nd ed. Frankfurt/Main: Degussa. Endter, R (1958), Chem .-Ing. -Tech. 30(5), 305. Fertig, W. A., Johnson, D. C , Dehong, L. E., McCallum, R. W, Maple, M. B., Bathas, B. T. (1977), Phys. Rev. Lett. 38(7), 987. Fiberg, M., Edwards, I. R. (1978), Report No. 1996, Nat. Inst. Metall. Randburg, Rep. of South Africa. Flanagan, T. B., Sakamoto, Y. (1993), Plat. Met. Rev. 37, 26. Fleischer, A., Lander, X X (1971), Zinc-Silver-Oxide Batteries. New York: Wiley. Fleischer, R. L. (1963), Acta Met. 11, 203. Fleischer, R. L. (1964), in: The Strengthening of Metals: Peckner, D. (Ed.). London: Reinhold, p. 93. Flukiger, R., Paoli, A., Roggen, R. Yvou, K., Muller, X (1972), Solid State Commun. 11(1), 61. Frembs, D. (1995), in: Edelmetalltaschenbuch, 2nd ed. Frankfurt/Main: Degussa. Friend, C. M. (1993), 5c/. Am. April, 42. Funston, E. S., McGurty, X A. (1961), Plat. Met. Rev. 5, 56. Furuya, K., Kallmann, S. (1991), in: Precious Metals: Benner, L. S. (Ed.). Allentown, PA: Int. Prec. Met. Inst. (IPMI). Gafner, G. (1989), Gold Bull. 22(4), 112. Gau, G., X-S. (1990), JOM, Feb., 35. Gerner, R., Jung, V., Kleinwiichter, L, Lang, J. (1991), in: Metall Forschung und Entwicklung, ed. Degussa Ressort Metall, Degussa, Frankfurt/ Main. p. 40.

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Girgis, K. (1983), in: Physical Metallurgy, 3rd ed.: Cahn, R. W, Haasen, P. (Eds.). Amsterdam: North Holland, Chap. 5. Glaser, H. J. (1989), Glastech. Ber. 62(3), 93. Green, D. (1972), /. Inst. Met. 10, 295. Gurney, P. D. (1993), Plat. Met. Rev. 37(3), 130. Haasen, P. (1979), in: Dislocations in Solids; Nabarro, F. R. N. (Ed.). Amsterdam: North-Holland, Chap. 15a. Haasen, P. (1983), in: Physical Metallurgy, 3rd ed.: Cahn, R. W., Haasen, P. (Eds.). Amsterdam: North Holland, Chap. 21. Haasen, P. (1984), in: Physikalische Metallkunde: Berlin: Springer. 1: p. 119; 2: p. 290; 3: p. 289; 4: p. 290. Hagen, E., Rubens, H. (1903), Ann Phys. 11, 873. Harriman, A. (1983), Plat. Met. Rev. 27(3), 102. Harriman, A. (1991), Plat. Met. Rev. 35(1), 22. Hass, G., Hunter, W. R. (1974), in: Proc. 9th International Congress of the International Commission for Optics, Santa Monica, CA, Oct. 1972: Thompson, B. X, Shannon, R. R. (Eds.). Washington, DC: Nat. Acad. Sci., pp. 525-553. Herzig, H. (1983), Plat. Met. Rev. 27(3), 108. Herzig, H., Spencer, R. S. (1982), Appl. Opt. 21(1), 15. Hirabayashi, M. (1991), in: Precious Metals: Brenner, L. S. (Ed.). Allentown, PA: IPMI, p. 105. Hiraga, K., Shindo, D., Hirabayashi, M., Terasaki, O., Watanabe, D. (1980), Acta Crystallogr. B36, 2550. Hoare, J. P. (1969), J. Electrochem. Soc. 116, 1131. Hume-Rothery, W. (1956), in: The Structure of Metals and Alloys: London: The Institute of Metals, 1: p. 127; 2: p. 109; 3: p. 205; 4: p. 229; 5: p. 193; 6: p. 132; 7: p. 110; 8: p. 197; 9: p. 192. Hume-Rothery, W. (1967), Prog. Mater. Sci. 13(5). Hume-Rothery, W, Reynolds, P. W. (1937), Proc. R. Soc. A160, 282. Hume-Rothery, W, Mabbot, G. W., Channel-Evans, K. M. (1934), Phil. Trans. R. Soc. A233, 1. Hummel, R. E. (1971), in: Optische Eigenschaften von Metallen und Legierungen: ed. W. Koster. Berlin: Springer, 1: p. 161; 2: p. 167; 3: p. 177; 4: p. 178. Humpston, G., Jacobson, D. M. (1992), Gold Bull. 25(4), 132. Humpston, G., Jacobson, D. M., Sangha, S. P. S. (1993), Gold Bull. 26(3), 90. Hunt, L. B. (1979), Gold Bull. 12, 116. Hunter, J. B. (1960), Plat. Met. Rev. 4(4), 130. Ishiyana, M. (1973), Ind. Rare Met. 52, 116. Iwasawa, Y. (1991), in: Precious Metals, Benner, L. S. (Ed.). Allentown, PA: p. 247. Johnson Matthey, (1992), Platinum. Johnson Matthey, (1993), Platinum. Johnson Matthey (1994), Platinum. Jonsson, S. (1995), in: Edelmetalltaschenbuch, 2nd ed. Frankfurt/Main: Degussa. Kempf, B., Hausselt, J. (1992), Interdiscip. Sci. Rev. 17, 251.

Khan, H. R. (1984), Gold Bull. 17(3), 94. Khare, S., Gontia, N., Nema, S. (1994), Plat. Met. Rev. 38, 175. Kidani, Y. (1991), in: Benner, L. S., Suzuki, T, Meguro, K., Tanaka, S. (eds.), Precious Metals Science and Technology, Allentown PA, Int. Prec. Met. Inst. (IPMI), p. 351. Kittel, C. (1986) Introduction to Solid State Physics, 6th ed. New York: Wiley. Knapton, A. G. (1977), Plat. Met. Rev. 21(2), 44. Kobayashi, K. (1983), in: Exotic Metals. Agune, p. 253. Kodas, T T, Hampden-Smith, M. J. (1994), The Chemistry of Metal CVD. Weinheim: VCH. Kondo, J. (1964), Prog. Theor. Phys. 32, 37; ibid. 35, 372. Kondo, J. (1969), Solid State Phys. 23, 184. Kopp, J. (1976), Gold Bull. 9, 55. Krappitz, H., Starz, K. A., Weise, W. (1991), in: Me tall Forschung und Entwicklung, ed. Degussa Ressort Metall, Degussa, Frankfurt/Main, p. 225. Kroogmann, K., Hausen, H. D. (1968), Z. Anorg. Allg. Chem. 358, 67. Kussman, A., Jessen, K. (1962), /. Phys. Soc. Jpn.

17 (Bl), 136. Kussman, A., Von Rittberg, G. (1950), Ann. de Phys. 7, 173. Labusch, R. (1970), Phys. Status Solidi 41, 659. Labusch, R. (1981), Cz. J. Phys. B31, 165. Lefort, T. E. (1931/35), French Patent 729 952, U.S. Patent 1998 878. Lewis, F. A. (1967), The Palladium-Hydrogen System. New York: Academic. Lewis, F. A. (1982), Plat. Met. Rev. 26, 121. Lewis, F. A. (1994), Plat. Met. Rev. 38(3), 112. Linde, J. O. (1932), Ann. Phys. (Leipzig) 14(5), 353; ibid. 15, 219. Lingen, D. (1983), Handbook of Batteries and Fuel Cells. New York: McGraw-Hill. Littnanski, H. (1965), Mitt. Forschungsges. Blechverarb., No. 4. Loebich, O. (1971), Doctoral Thesis, I E Aachen. Madix, R. J. (1986), Science 233, 1159. Maeno, Y, Hashimoto, H., Yoshida, K., Nishizaki, S., Fujita, T, Bednarz, J. G., Lichtenberg, F. (1994), Nature 372, 532. Mahvan, N., Zeltser, A. M., Lambeth, D. N., Laughlin, D. E., Kryder, M. H. (1990), IEEE Trans. Magn. 26, 2277. Maier, C , Bilger, G. (1994), Plat. Met. Rev. 38,175. Maple, M. P., Hamaker, H. C , Woolf, L. D., Mackay, H. B., Fisk, Z., Odoni, W, Ott, H. R. (1980), in: Crystalline Electric Field and Structural Effect in f-Electron Systems, Crow, J. E. (Ed.). New York: Plenum, p. 533. Massalski, T. B. (1983), in: Physical Metallurgy, 3rd ed.: Cahn, R. W, Haasen, P. (Eds.). Amsterdam: North-Holland, Chap. 4. Massalski, T. B., Pearson, W. B., Bennett, L. H., Chang, Y A. (Eds.) (1986), Noble Metal Alloys:

9.7 References

Phase Diagram, Alloy Phase Stability, Thermodynamic Aspects and Special Features, Warrendale, PA: The Metallurgical Society, USA. Matucha, K. H. (1982), Bericht Nr. 72 des Metalllaboratoriums der Metallgesellschaft, Frankfurt/ Main. McKenzie, D. R. (1978), Gold Bull 11, 49. Mills, A. (1991), Chemistry of the Platinum-Group Metals: in: Hartley, F. R. (Ed.). Amsterdam, Elsevier, pp. 302-337. Minot, M. X, Perlstein, J. H. (1971), Phys. Rev. Lett. 26, 371. Mizuhara, H. (1987), Met. Prog. 131(2), 53. Mohan, A. (1976), Gold Bull. 9, 66. Moore, G. A. (1939), Trans. Electrochem. Soc. 75, 237. Motoo, S. (1991), in: Benner, L. S., Suzuki, T., Meguro, K., Tanaka, S. (eds.), Precious Metals Science and Technology, Allentown PA, Int. Prec. Met. Inst. (IPMI), p. 342. Nadler, M. R., Kempter, C. P. (1960), J. Phys. Chem. 64, 1468. Nelin, G. (1971) Phys. Stat. Solidi 45, 527. Nicholas, M., Poole, D. M. (1966), Trans. Metall. Soc. AIME236, 1535. Nicholson, E. D. (1979), Gold Bull 12(4), 161. Normandeau, G. (1992), Gold Bull. 25(3), 94. Obuchi, A., Naito, S., Onishi, X, Tamaru, K. (1982), Surf. Sci. 122, 235. Obuchi, A., Naito, S., Onishi, T., Tamaru, K. (1983), Surf Sci. 130, 29. Olsen, D. R., Berg, H. M. (1979), IEEE Trans. CHMT-2, 257. Osborne, J. A., Wilkinson, G., Young, X F. (1965), /. Chem. Soc, Chem. Commun. 131. Ostwald, W. (1894), Z. Phys. Chem. 15, 705. Ostwald, W, Brauer, E. (1931), in: Festschrift zum 50jdhrigen Bestehen der Platinschmelze: ed. Huben, H., Hanau: G. Siebert, pp. 240-256. Ott, D., Raub, C. X (1980/81/82), Z. Metall 34, 629; 55, 543, 1005; 36, 150. Panfilov, P., Yermakov, A. (1994), Plat. Met. Rev. 38(1), 12. Peierls, R. E. (1955), Quantum Theory of Solids. Oxford, Clarendon, p. 108. Peyrone, M. (1845), Ann. 51, 15. Pollitzer, E. L. (1972), Plat. Met. Rev. 16(2), 42. Poniatowsky, M. (1995), in: Edelmetalltaschenbuch, 2nd ed. Frankfurt/Main: Degussa. Poniatowsky, M., Clasing, M. (1968), Z. Metallkd. 59, 165. Prescher, G. (1995), in: Edelmetalltaschenbuch, 2nd ed. Frankfurt/Main: Degussa. Prescher, G., Lehmann, Y. (1995), in: Edelmetalltaschenbuch, 2nd ed. Frankfurt/Main: Degussa. Prescher, G., Seybold, K. (1995), in: Edelmetalltaschenbuch, 2nd ed. Frankfurt/Main: Degussa. Pugh, S. F. (1954), Phil Mag. 45, 823. Raub, C. X (1984), Plat. Met. Rev. 28(2), 63.

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Raub, E. (1940), in: Die Edelmetalle und ihre Legierungen: ed. W. Koster, Berlin: Springer, 1: p. 274, 2: pp. 4, 45; 3: pp. 12, 49. Raub, E. (1959), J. Less Common Met. 1, 3. Raynor, G. V. (1976), Gold Bulletin 9. 1: p. 17, 2: p. 50, 3: p. 52. 4: p. 16; Regenet, P. X, Stiiwe, H. P. (1963), Z. Metallkd. 54, 273. Renner, H. (1990), in: Ullmann's Encycl, Vol. A12, 1: p. 524, 2: p. 525, 3: p. 521, 4: p. 517. Renner, H. (1992), in: Ullmann's Encycl, Vol. A21, p. 114. Renner, H. (1993), in: Ullmann's Encycl, Vol. A24, l : p . 156,2: p. 144, 3: p. 125. Riabkina, M., Gal-Or, L. (1984), Gold Bull 17(2), 62. Richardson, F. D. (1974), in: Physical Chemistry of Melts in Metallurgy: London: Academic Vol. 1. Rosenberg, B., van Kamp, L., Krigus, T. (1965), Nature 203, 698. Rosenberg, B., Trosko, X E., Mansour, V. H. (1969), Nature 222, 385. Sato, H., Toth, R. S. (1964), in: Alloy, Behav. a. Effects in Cone. Sol. Solutn., Massalski, T. B. (Ed.). New York: Gordon & Breach, p. 295. Savage, R., Tracey, A. (1971), Mod. Devel Powder Met. 5, 273. Savitskii, E. M., Prince, A. (1989), in: Handbook of Precious Metals: New York: Hemisphere. Schlott, M. (1995), in: Edelmetalltaschenbuch, 2nd ed. Frankfurt/Main: Degussa. Schmid, E., Wassermann, G. (1925), Z. Metallkd. 19, 386. Schroeder, K. (1995), in: Edelmetalltaschenbuch, 2nd ed. Frankfurt/Main: Degussa. Schryvers, D., van Landuit, X, van Tendeloo, G., Amelinckx, S. (1983), Phys. Status Solidi (a) 76, 575. Schuit, G. C. A., Van Reijen, L. L. (1958), Adv. Catal 10, 242. Schulze, G. E. R. (1974), in: Metallphysik. Springer New York, p. 118. Selman, G. L., Day, X G., Bourne, A. A. (1974), Plat. Met. Rev. 18(2), 46. Seibold (1995), in: Edelmetalltaschenbuch, 2nd ed. Frankfurt/Main: Degussa. Shul'ga, Y. M., Izakovich, E. N., Rubtsov, V I., Shklyaruk, B. F. (1993), Plat. Met. Rev. 37(2), 86. Sicking, G. (1970), Ph.D. Thesis, Minister, Germany. Simon, F. (1992), Gahanotechnik 83, 808; 1180. Simon, R, Zilske, A. (1995), in: Edelmetalltaschenbuch, 2nd ed. Frankfurt/Main: Degussa. Sistare, G. H. Jr., McDonald, A. S. (1979), in: Metals Handbook, 9th ed., Vol. 2, Metals Park, OH: ASM, p. 681. Smith, F. X (1975), Plat. Met. Rev. 19, 93. Smithells, C. X, Brandes, E. A. (1977), in: Metals Reference Book, 5th ed. London: Butterworth, p. 1019.

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Sperner, R, Hohmann, W. (1976), Plat. Met. Rev. 20, 12. Stanley-Wood, N. (1977), Powder Met. Int. 9, No. 3; Powder Metall. Rev. 10(3), 138. Starz, K. A. (1995), in: Edelmetalltaschenbuch, 2nd ed. Frankfurt/Main: Degussa. Stockel, D. (1974), Z. Metall. 28(7), 611. Stockel, D. (1979), Z. Werkstofftechnik 10, 1: p. 230; 2: p. 238. StraBer, K. (1990) Ber. Bunsenges. Phys. Chem. 94, 1000-1005. Tachiki, M., Teramoto, K. (1966), / Phys. Chem. Solids 27, 335. Takeuchi, S. (1972), in: The Properties of Liquid Metals - Proc. 2nd Int. Conf. Tokyo, 1972: London: Taylor and Francis. Tamaru, K., Naito, S. (1991), in: Precious Metals: Benner, L. S. (Ed.). Allentown, PA: IPMI, p. 262. Tanaka, T. (1991), in: Precious Metals: Benner, L. S. (Ed.). Allentown, PA: IPMI, 1: p. 335; 2: p. 336. Tanino, H. (1982), Kotaibutsuri 17, 728. Tanuma, S. (1991), in: Precious Metals: Benner, L. S. (Ed.). Allentown, PA: IPMI, p. 46. Taylor, K. C. (1993), Catal. Rev.-Sci. Eng. 35(4), 457. Tennant, S. (1804), Phil. Trans. R. Soc. 94, 411. Terasaki, O., Watanabe, D., Hiraga, K., Shindo, D., Hirabayashi, M. (1981), J. Appl. Crystallogr. 14, 392. Ullmann's Encyclopedia of Industrial Chemistry (1993), Vol. A24. Weinheim: VCH, pp. 107-163. Vannice, M. A. (1975), J. Catal 37, 449; 50, 228. Vannice, M. A. (1977), J. Catal. 50, 228. Volcker, A. (1995), in: Edelmetalltaschenbuch, 2nd ed. Frankfurt/Main: Degussa. Wassermann, G. Grewen, J. (1962), Texturen metallise her Werkstoffe, 2nd ed. Berlin: Springer. Watkins, D. M. (1998), Plat. Met. Rev. 22(4), 118. Watson, W. (1951), Phil. Trans. R. Soc. 46, 584. Weise, W. (1991), in: Metall Forschung und Entwicklung (1991), ed. Degussa Ressort Metall, Degussa, Frankfurt/Main, p. 177. Weller, D., Brandle, H., Gorman, G. L., Lin, C. J., Notarys, H. (1992), Appl. Phys. Lett. 61, 2276. Werner, A. (1893), Z. Anorg. Chem. 3, 267. Wicke, E., Nernst, G. H. (1964), Ber. Bunsenges. 68, 224. Wise, E. M. (1964), Gold, Van Nostrand, p. 72. Wise, E. M., Vines, R. F. (1979), in: Metals Handbook, 9th ed., Vol. 2. Metals Park, OH: ASTM, p. 699. Wollaston, W. H. (1804), Phil. Trans. R. Soc. 94, 419. Wollaston, W. H. (1805), Phil. Trans. R. Soc. 95, 316. Yamauchi, H., Yoshimatsu, H. A., Forouhi, A. R., de Fontaine, D. (1981), in: Precious Metals: McGachie, R. O., Bradley, A. G. (Eds.). Toronto: Pergamon, p. 241. Yasuda, K., Hisatsune, K. (1993), Gold Bull. 26(2), 50.

Yates, J. T, Jr. (1992), Chem. Eng. News, March 30, 22. Yogi, T, Tsang, C , Nguyen, T. A., Yu, K., Gorman, G. L., Castillo, G. (1990), IEEE Trans. Magn. 26, 2271. Yonemitsu, K. (1991), in: Benner, L. S., Suzuki, T., Meguro, K., Tanaka, S. (eds.), Precious Metals Science and Technology, Allentown PA, Int. Prec. Met. Inst. (IPMI), p. 25. Zimmermann, K. F. (1966), Degussa-Berichte aus Forschung und Praxis, Nr. 16, Hartloten. Zysk, E. D. (1979), in: Metals Handbook, 9th ed., Vol. 2. Metals Park, OH: ASM, p. 689.

General Reading Benner, L. S. (1991), Precious Metals, Science and Technology. Allentown, PA: International Precious Metals Institute. Bohm, H. (1968), Einfuhrung in die Metallkunde. Mannheim: Bibliographisches Institut. Darling, A. S. (1973), Some Properties and Applications in the Platinum Group Metals. London: Institute of Metals. Degussa (1967), Edelmetall-Taschenbuch. Frankfurt/ Main: Degussa. Degussa (1991), Metall Forschung und Entwicklung. Frankfurt/M.: Degussa, Ressort Metall. Degussa (1995), Edelmetall-Taschenbuch, 2nd ed. Heidelberg: Hiithig Verlag. Doduco Datenbuch (1977), Pforzheim: Doduco Kg. Elliot, R. P. (1965), Constitution of Binary Alloys, First Supplement. New York: McGraw-Hill. Fluck, E., Heumann, K. G. (1986), Periodic Table of the Elements. Weinheim: VCH. Gerald, V. (Ed.) (1993), Structure of Solids, Materials Science and Technology, Vol. 1. Weinheim: VCH. Haasen, P. (1984), Physikalische Metallkunde. Berlin: Springer. Hansen, M. Anderko, K. (1958), Constitution of Binary Alloys. New York: McGraw-Hill. Hering, E., Schulz, W (1988), Periodensystem der Elemente, Diisseldorf: VDI. Hume-Rothery, W, Raynor, G. V. (1956), The Structure of Metals and Alloys. London: The Institute of Metalls. Hummel, R. E. (1971), Optische Eigenschaften von Metallen und Legierungen. Berlin: Springer. Kittel, C. (1986), Introduction to Solid State Physics, 6th ed. New York: Wiley. Landolt-Bornstein (1967), 6th ed. Berlin: Springer, Bd. IV/a, pp. 66-75. Landolt-Bornstein (1985), New Series, Vol. III/15b; Berlin: Springer, pp. 222-301. Massalski, T. B. (1987), in: Binary Alloy Phase Diagrams, Vols. 1 & 2. Metals Park, OH: American Society for Metals.

9.7 References

Metals Handbook (1979), Ninth Ed., Vol. 2, Properties and Selection: Nonferrous Alloys and Pure Metals. Metals Park, OH: American Society for Metals, pp. 659, 671, 684, 689, 699. Petzow, G., Effenberg, G. (Eds.) (1988), Ternary Alloys, Vols. 1 and 2. Weinheim: VCH. Raub, E. (1940), Die Edelmetalle und ihre Legierungen. Berlin: Springer. Raynor, G. V. (1976), Gold Bull I: p. 12-18, II: pp. 50-55. Savitskii, E. M., Prince, A. (1989), Handbook of Precious Metals. New York: Hemisphere. Shunk, F. A. (1969), Constitution of Binary Alloys, 2nd Suppl. New York: McGraw-Hill.

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Smithells, C. X, Brandes, E. A. (1977), in: Metals Reference Book, 5th. ed. London: Butterworth. Ullmann's Encyclopedia of Industrial Chemistry (1990), Vol. A12. Weinheim: VCH, pp. 499-533. Ullmann's Encyclopedia of Industrial Chemistry (1992), Vol. A21. Weinheim: VCH, pp. 75-131. Ullmann's Encyclopedia of Industrial Chemistry (1993), Vol. A24. Weinheim: VCH, pp. 107-163. Villars, P., Calvert, L. D. (1985), Pearson's Handbook of Crystallographic Data for Intermetallic Phases, Vols. 2 & 3. Metals Park, OH: American Society for Metals.

10 Refractory Metals and Their Alloys Erwin Pink

Erich-Schmid-Institut fur Festkorperphysik der Osterreichischen Akademie der Wissenschaften, Leoben, Austria Ralf Eck

Metallwerk Plansee Ges.m.b.H., Reutte, Austria

List of Symbols and Abbreviations List of Material Designations 10.1 Introduction 10.2 Production Methods 10.2.1 General Remarks and Production Equipment 10.2.1.1 Compacting 10.2.1.2 Sintering 10.2.1.3 Melting 10.2.1.4 Forming 10.2.1.5 Heat Treatment 10.2.2 Production of Individual Metals 10.2.2.1 Vanadium 10.2.2.2 Niobium and Tantalum 10.2.2.3 Chromium 10.2.2.4 Molybdenum and Tungsten 10.2.2.5 Rhenium 10.2.3 Recycling 10.3 Characteristics of Pure Refractory Metals 10.3.1 Physical Properties 10.3.2 Chemical Properties 10.3.3 Recrystallization Behavior 10.3.4 Important Mechanical Properties 10.4 Structure and Properties of Refractory Alloys 10.4.1 Substitutional Alloys 10.4.2 Doped Tungsten and Molybdenum 10.4.3 Dispersion-Strengthened Alloys 10.4.4 Complex Alloys 10.4.5 Refractory Alloys for Electrical Contacts 10.4.6 Tungsten Heavy Alloys 10.4.7 Composites Reinforced with Refractory-Metal Fibers 10.4.8 Refractory-Metal Cermets 10.4.9 Amorphous Refractory Alloys Materials Science and Technology Copyright © WILEY-VCH Verlag GmbH & Co KGaA. Allrightsreserved.

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10.5 10.5.1 10.5.2 10.5.3 10.5.3.1 10.5.3.2 10.5.3.3 10.5.3.4 10.5.3.5 10.6 10.6.1 10.6.1.1 10.6.1.2 10.6.1.3 10.6.2 10.6.2.1 10.6.2.2 10.6.2.3 10.6.2.4 10.7 10.7.1 10.7.2 10.7.3 10.7.3.1 10.7.3.2 10.7.4 10.7.5 10.7.5.1 10.7.5.2 10.7.5.3 10.7.5.4 10.7.5.5 10.7.5.6 10.7.5.7 10.7.6 10.7.6.1 10.7.6.2 10.7.6.3 10.7.7 10.7.7.1 10.7.7.2 10.7.7.3 10.8 10.9

10 Refractory Metals and Their Alloys

Applications of Refractory Metals and Alloys Statistics General Applications Requirements for Special Applications The Quest for High Purity Porous Metals Refractory Alloys for Thermocouples Refractory Metals for Superconductors Requirements for Use in Thermonuclear Reactors Further Processing Fastening and Joining Mechanical Fastening Welding Brazing Coatings Pack and Slurry Cementation Electrolytic Coating Flame and Plasma Spraying Chemical and Physical Vapor Deposition Fundamental Mechanisms Activated and Liquid-Phase Sintering Interstitial Solubilities Introducing Dispersed Second Phases Carbide Particles Other Particles Recrystallization in Alloys Micromechanisms of Deformation and Strengthening Solid-Solution Strengthening Dynamic Strain Aging The Effect of Dispersed Second Phases Thermomechanical Treatments Grain-Size and Grain-Shape Strengthening Rhenium-Alloyed Refractory Metals: Deformation Twinning The Deformation of Pseudo-Alloys and Tungsten Heavy Alloys Micromechanisms of Fracture The Ductile-Brittle Transition Temperature (DBTT) High-Temperature Fracture of Non-Sag Structures Fracture of Tungsten Heavy Alloys Structure-Related Anelastic Properties Snoek-Type Damping Damping from "Particle" Formation and Dissociation Grain-Boundary Damping Recent Developments and Future Prospects References

616 616 617 619 619 619 619 619 620 621 621 621 621 622 622 623 623 624 624 625 625 625 626 626 627 627 628 629 629 630 631 631 632 632 633 633 635 635 636 636 636 637 637 638

List of Symbols and Abbreviations

List of Symbols and Abbreviations A C E

f I n S So T T W Z

a

fracture elongation contiguity Young's modulus grain-aspect ratio fracture toughness length; grain length stress exponent cross section after working original cross section of sinter ingot temperature melting temperature grain width constant

G

coefficient of linear thermal expansion strain rate shear modulus stress

APS b.c.c. CVD DBTT EB f.c.c. HB, HV h.c.p. H.T.D.C. L.T.D.C. Me P/M PVD TIG UTS VPS YS

plasma spraying in air body-centered cubic chemical vapor deposition ductile-brittle transition temperature electron beam face-centered cubic Brinell, Vickers hardness (HV10: 100 N load) hexagonal close-packed high-temperature dislocation creep low-temperature dislocation creep metal powder metallurgy, powder metallurgical physical vapor deposition tungsten inert-gas ultimate tensile strength plasma spraying in vacuum vield stress

8

n

List of Material Designations AKS KS MHC ODS

doped with aluminum-potassium silicates doped with potassium silicates molybdenum alloy containing Hf and C oxide dispersion-strengthened

591

592

PMS SMS TZM ZHM

10 Refractory Metals and Their Alloys

Chevrel phase based on Pb, Mo, and S Chevrel phase based on Sn, Mo, and S molybdenum alloy containing Ti and Zr (and C) molybdenum alloy containing Zr and Hf (and C)

Alloy characterization: e.g., W-5Mo denotes a W alloy with 5 wt.% Mo; W-(l-5)Mo stands for a range 1-5 wt.% Mo; W-(l;5)Mo are alloys with 1 and 5 wt.% Mo.

10.2 Production Methods

10.1 Introduction The metals of groups Va and Via of the periodic system of elements are commonly called "refractory metals". They are vanadium, niobium and tantalum, and chromium, molybdenum and tungsten. They all have high melting points, culminating at 3410 °C for tungsten, which is why their mechanical properties can be exploited to much higher temperatures than those of the common metals and alloys. Following another definition, refractory metals are those with a melting point above 2000 °C. This would comprise ten metals, some of which are described in Chaps. 8 and 9 of this Volume. Rhenium of group Vila will be included here as it does not fit into any other classification. In recent years it has gained importance as a refractory metal for exceptional cases and as an alloying element for group Via metals. Several books (e.g., Northcott, 1956; Rostoker, 1958; Kieffer and Braun, 1963; Tietz and Wilson, 1965; Rieck, 1967; Sully and Brandes, 1967; Wilkinson, 1969; Kieffer et al, 1971; Trefilov and Moiseev, 1978; Benesovsky, 1979; Yih and Wang, 1979; Pink and Bartha, 1989), reference works, conference proceedings and numerous original papers in periodicals have been published on the technology of refractory metals and on metal-physics topics1. The goal of the present review is to briefly treat general production methods in relation to resulting properties, and to emphasize those relations between structure and properties which are observed mainly for this group of metals. The discussion will in general be restricted to 1 For the sake of brevity, recent review articles and books rather than original papers are quoted in the reference list. Publications dealing with general principles are not included.

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those metals and alloys which are commercially available.

10.2 Production Methods 10.2.1 General Remarks and Production Equipment

Today the majority of refractory metals are produced by means of powder-metallurgical methods (see Chap. 4 by Arunachalam and Sundaresan, in Vol. 15 of this Series). 10.2.1.1 Compacting Powders of all commercially fabricated refractory metals and alloys are consolidated before sintering to reach densities averaging 60-65% (75% at most) of the theoretical, mainly by one of two methods: • Pressing of powders in rigid dies on hydraulic presses using pressures of about 3000 bar (300 MPa). Dimensional precision is high using mechanical presses, but hydraulic presses allow better control of the ram movement. Molybdenum blanks can be machined without sintering, but tungsten has to be pre-sintered. • Isostatic pressing of powders in flexible containers. For very large parts, isostatic wet-bag pressing is the most common method, and leads to homogeneous densities of 60-65%. Other methods, such as extrusion, powder rolling, vibratory or explosive consolidation, and slip casting have only marginal importance. 10.2.1.2 Sintering Pre-sintering of billets compacted from powder facilitates handling and machining. The primary goal of sintering is densification (to at least 90 % of the theoretical

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10 Refractory Metals and Their Alloys

density), which is required to make further working operations successful. In addition to ensuring high density, purification occurs during sintering. Typical particle sizes of metal powders are 1-10 jim; the grain sizes after sintering are 10-30 \xm. Sintering is carried out in different ways: • Self-resistance heating or "direct sintering" is used extensively for tungsten because of the extremely high temperatures required for practical sintering. Rods are heated to 3000 °C in a hydrogen atmosphere. Tantalum and its alloys are sintered at 2600 °C in vacuum. • Horizontal furnaces with resistanceheating elements are frequently used for hydrogen sintering of molybdenum and tungsten. A temperature of 1800°C is a practical limit for furnaces that use aluminum-oxide refractories. Ceramic-free, high-temperature furnaces with heating elements of tungsten sheet and heat shielding made from molybdenum alloys can operate at 2700 °C. • Induction radiant heating can also be used. A typical furnace consists of a chamber that contains a cylindrical susceptor ring of tungsten or molybdenum. The susceptor is inductively heated and the compact is heated by radiation from the susceptor up to 2400 °C. Among the many factors which influence sintering, the following are of interest for the refractory metals: • Time and temperature are fundamental determinants of the sintering kinetics. Mass transport by diffusion is very slow for refractory metals at sintering temperatures that are economically or technically feasible. • Purity. The impurities in the compact greatly affect the sintering of molybdenum and tungsten, as both metals are sintered at high temperatures, at which most impu-

rities are molten, and many have high vapor pressures. The entrapment of gases prevents the achievement of high density. Carbon causes carburization and inhibits densification. Metallic impurity elements may act as sintering activators (see Sec. 10.7.1) and may leave large voids where the contaminant was originally present. • Thermal gradients and heating rates are particularly important when large furnace loads are sintered. Because sintering rates are so sensitive to temperature, thermal gradients greatly affect the density attained in a given time. At 1800-2500 °C, radiation is the predominant mode of heating and the billet surfaces reach sintering temperatures much more rapidly than the billet interiors. If not controlled, a thin high-density shell forms, which effectively restricts further densification in the interior. 10.2.1.3 Melting The advantage of melting is the higher degree of homogeneity obtained in the alloys. However, the coarse grain structure that results from melting makes further processing difficult. In order to break up the large crystals in the ingots of molybdenum, they must be extruded, which raises the already high costs still further. For niobium, tantalum and some alloys, electronbeam and plasma melting are the most important methods. 10.2.1.4 Forming In contrast to standard metal-forming operations there are two important details to be observed when dealing with sintered refractory metals. Initial forming has to take into account the varying number of pores which are present in the sintered condition. Reduction is necessary to such

10.2 Production Methods

an extent as to achieve fully dense products and to eliminate any effects of pores, e.g., trapped gases or crack initiation. Furthermore, if intermediate heating is necessary, it must be performed in a protective atmosphere, i.e., in vacuum or oxygen-free, inert gases, or in (cheaper) hydrogen-flushed furnaces (in the case of Via metals) to avoid oxidation or sublimation of oxides. For the same purpose, the sintered ingots can be canned in mild steel. Ingots of group Va metals (which are softer and more ductile) are cold rolled (at a maximum temperature of 300-400 °C) to avoid oxidation. Plastic forming of molybdenum, tungsten and their alloys is more difficult. The initial heating temperature necessary for molybdenum is 13001400°C, and 1500-1600°C for tungsten. During mechanical forming, the intermediate heating temperatures are successively lowered. After severe deformation (to produce thin wires, sheets or foils), molybdenum and its alloys can be processed at room temperature, whereas tungsten and its alloys are rolled and drawn at temperatures of 500-200 °C without a protective atmosphere. Molybdenum alloys, tungsten, and tungsten alloys do not normally recrystallize during processing. Only pure molybdenum may recrystallize during heating at standard temperatures. All of the forming processes which are common in other metal-producing industries are equally suited for refractory metals: • Molybdenum, tungsten, and their alloys, for example, weighing up to 600 kg are hot- and cold rolled to 20 mm thick sheet. Rods of 25-50 mm diameter are rolled on grooved rolls. Foils have a minimum thickness of 15-30 jim. • Forging in dies is standard practice for simple, radially symmetrical parts with maximum diameters of about 200 mm.

595

TZM ingots weighing 4000 kg have been upset-forged between flat dies. • Rotary swaging allows the production of rods down to a few millimeters in diameter and several meters in length. • By extruding, superior strength in comparison to rolling or forging is obtained in longitudinal and transverse directions for rods and thick-walled tubes. • Flow turning is used for the fabrication of rotationally symmetrical shapes, starting from discs to form cones and cylinders. The disc, which is heated with a torch, is continuously reduced in thickness. The largest discs, flow-turned to 5-8 mm in thickness, start at 1.6 m in diameter for molybdenum or TZM. • Differing from flow turning, spinning is a process to form cones and cylinders from a round plate of 1 - 5 mm in thickness and 1.5 m in diameter, without a substantial further decrease in the thickness for molybdenum and tungsten. • Deep drawing is an alternative to spinning. The starting material should be a cross-rolled sheet to avoid extensive earing. • For wire fabrication, sintered square or round rods are first rolled on grooved rolls, and then they are further reduced in diameter by drawing through dies. Thin wires of 5 jim in diameter can be drawn using drawing machines that serve ten or more dies with one traction. While extrusion and spinning are applied for producing tubes of short lengths, the drawing process is also suited for the fabrication of long, seamless tubes, starting from an extruded tube or a drilled rod. 10.2.1.5 Heat Treatment The goal of a heat treatment of refractory metals is to minimize stresses (to counteract work hardening), and to increase the

596

10 Refractory Metals and Their Alloys

elongation through beginning recrystallization. In alloys, the precipitation of phases which strengthen or embrittle may be controlled by selected time-temperature cycles. Annealing of worked metals below their recrystallization temperature may relieve residual stresses and improve the ductility to a certain extent. Internal cracking and lamination of highly cold-worked wires and sheets can be avoided by intermediate, partial or complete recrystallization. Heat treatments between certain stages of forming and afterwards are also performed to control the final mechanical properties. Heat treatments are performed in vacuum or an oxygen-free atmosphere (Va metals), or under hydrogen (Via metals). The advantages of heating in vacuum are reproducible conditions and the cleanest atmosphere. One of the problems of heat treating alloys is the oxidation of alloying elements such as titanium, zirconium, and hafnium, and the depletion of carbon. In contrast to unalloyed molybdenum and tungsten, such alloys must not be annealed in nitrogen-containing atmospheres because nitrides formed on the surfaces have an embrittling effect. 10.2.2 Production of Individual Metals 10.2.2.1 Vanadium

The commercial production of vanadium began comparatively late in the 1960s (Anonymous, 1988 a). Vanadium powder is obtained by an aluminothermic reduction. Electron-beam melting is applied to produce high-purity ingots weighing up to 1.5 tons (1.5 x 103 kg). All of the standard dimensions of semi-finished products can be fabricated. For wire and thin-tube drawing it is desirable to cover the metal with a lubricating oxide layer to prevent it from sticking

to the tools. This is done by electro-chemical oxidizing, but formation appears to be difficult. An additional problem when vanadium alloys are drawn at higher temperatures is that the V 2 O 5 thus formed may be molten (its melting point is 660 °C), and since it is highly corrosive, it may attack the tools. 10.2.2.2 Niobium and Tantalum Tantalum powder, or a small amount of niobium powder is processed further by means of powder metallurgy (P/M). With self-resistance sintering, temperatures of 2500-2700 °C are reached, but the best possible vacuum is then determined by the high metal vapor pressure. Heavy niobium and tantalum products are usually fabricated using vacuum arccasting or, preferably, electron-beam melting techniques; their impurity levels are shown in Table 10-1. Vacuum melting purifies the starting block. The formation of homogeneous alloys is facilitated by melting techniques. Although the disadvantage of melting, which is the resulting coarse grain size, can be controlled, the fine grain size that can be achieved using sintering is never reached. As a prerequisite for wire or thin-tube drawing, the surfaces are covered with a Ta 2 O 5 layer, which is state of the art. 10.2.2.3 Chromium Various melting techniques were previously used to produce the purest possible chromium in the laboratory. The main problem was to reduce the brittleness of the chromium near room temperature. Disregarding the cost of production, it was verified that remelted chromium starting from "highesf'-purity iodide chromium can be ductile when purities above 99.998% were achieved (Heraeus, 1975).

597

10.2 Production Methods

Table 10-1. Typical impurity concentrations in commercial melted and P/M refractory metals (in jag/g). Impurity

Refractory metal Va

Nb a

Taa

Cr

Mo

KS-Mo

W

AKS-W

H 0 N C

<5 250 30 30

<1 20 <5 15

<1 300-500 <5 15

<1 <5 <5 15

<1 10 <5 15

<30 <20

3 15 <5 10 _ <5

3 80 30 40

P

<1 30 20 30 _ <5

<10 10

<5 <5

20 <2

20 <2

<50 <5 -

<5 <5 <20

10 <5 5

<10 <5 <10

10 50 <10

(99.3)b <10 <10 300 150 150 200 <10

6000 (99.99)b <10 <10 <20 <5 <5 <5

<5 <5 50 _ — (99.97) b 60 100 10 60 <5

<5 <5 <10 <20 (99.98) b 5 15 20

<5 <5 <10 40 (99.90)b 5 20 <10

s Al K Si Nb Ta Cr Mo W Ni Fe Pb a

300 400 60 <20 80 <20 <100

Commercial melted;

b

<5 <5 <10 (99.97)b 100 <5 30 <5

5 <5 10 150-300 500-900 <5 <5 <10 (99.84)b 100 5 50 <5

purity (wt.%).

P/M was also tried, but suitable production of pure powder, adequate processing, and the fabrication of a pore-free final product have only recently led to satisfactory results. A new approach using P/M methods was launched to economically produce 99.97 % pure chromium by means of comminution of electrolytical chromium, powder preparation, compaction, sintering, and plastic forming (Eck et al., 1990). Typical impurity concentrations are given in Table 10-1. Of these, nitrogen is considered the most difficult element to control during production, and is the most embrittling element.

tion from ore to finished products is shown in Fig. 10-1. Although sintering is the preferred way to produce molybdenum and tungsten, arc and electron-beam melting are also industrial practice for molybdenum and its alloys, although only an estimated 5% or less are produced in this way. All other melting methods are unsuccessful. The purities of molybdenum and tungsten products and their impurity contents are shown in Table 10-1. Besides low levels of hydrogen, oxygen, nitrogen, and carbon, low levels of phosphorus and iron are essential for processing sintered blanks.

10.2.2.4 Molybdenum and Tungsten

10.2.2.5 Rhenium

As an example of the many chemical reactions necessary to produce a metal powder, a flow sheet for tungsten produc-

P/M results in excellent purity (99.96%) and high material yield. After sintering at temperatures as used-for tungsten, densi-

598

10 Refractory Metals and Their Alloys

TUNGSTEN ORE

VERY PURE SCHEELITE

WOLFRAMITE AND SCHEELITE CONC.

{CONCENTRATES

SCHEELITE CONC. Fusion with soda

HC1Reduction (metallo- or carbothermic)

WOLFRAMITE CONC. Fusion with NaOH or KOH I

I SODIUM WOLFRAMATE]

CaCl, -HC1 SYNTHET. SCHEELITE

n«*-H,PO4 COLORS PIGMENTS CHEMICALS

TUNGSTEN ACID} <•—NH,OH

FERROTUNGSTEN|

AMMONIUMPARATUNGSTATE TUNGSTEN-ALLOYED STEEL Magnets, tool steels, hightemperature construction steels, high-speed steels

Calcined to WO, Reduction with C

NON-FERROUS ALLOYS, COMPOSITE MATERIALS High-temperature highstrength alloys, stellite, heavy alloys, electrical contacts TUNGSTEN-CARBIDE HARD METAL

Sintering -

Reduction in H2

I

TUNGSTEN POWDER Sintering Forming

TUNGSTEN POWDER

CARBIDE POWDER

ties of the sintered compacts reach 80-90 % of the theoretical. Electron-beam melting of rhenium is feasible but uneconomical. Plastic forming of pure rhenium is difficult for several reasons (Carlen and Bryskin, 1992): it work hardens heavily, so that recrystallization anneals become necessary after little cold rolling, and it can hardly be worked above room temperature because it embrittles due to the melting of a rhenium oxide, which melts at 300 °C and penetrates the grain boundaries. One way to avoid the inherent and the impurity-induced embrittlement (critical in swaging and drawing) is extruding, preferentially done in combination with cladding in tantalum. Under such conditions, extruding at temperatures as high as 2000 °C can be successful (Witzke and Raffo, 1971), but this is not done in practice.

COMMERCIAL TUNGSTEN Sheets, rods, wire, forgings

Figure 10-1. Flowsheet for the production of tungsten from ore to final product (Gocht, 1974).

10.2.3 Recycling Most of the scrap collected from fabrication or from used parts or constructions is converted by similar chemical methods as used for powder production from the ore. High-purity scrap which can be collected during the P/M production is recycled the easiest way: as high-purity alloying material for superalloys and steels. If oxygen or carbon are detrimental, then scrap has to be pickled in molten sodium hydroxide at 400 °C. Cleaned and recrystallized scrap can also be recycled by comminution milling at temperatures below — 40 °C, using liquid nitrogen as the coolant. Pick-up of iron from the mill can be minimized by lining the inside of the mill with a sheet of the metal to be ground. Tungsten powder of

10.3 Characteristics of Pure Refractory Metals

such origin is further processed to tungsten heavy metal, whereas molybdenum powder can be used for plasma spraying. Another process, with uses the cooling effect of the adiabatic expansion of nitrogen, is called the "coldstream process", and can be used for tungsten. Pure tungsten-copper and tungsten heavy metal can be recycled by an oxidation-reduction process to an alloy powder which can be processed directly. Vacuum distillation could be used to recover infiltrated tungsten, but is expensive. Tantalum scrap is, when contaminated, pretreated mechanically and chemically, and arc-cast or electron-beam melted. It is essential that rhenium, a "precious" metal, be recovered. Chemical processing is an established procedure, as is sublimation of the oxide.

10.3 Characteristics of Pure Refractory Metals 10.3.1 Physical Properties The metals of group Via have the highest melting points within their period. Tungsten as the highest-melting metal is surpassed only by carbon, and by the monocarbides of zirconium, hafnium, niobium and tantalum. All of the metals of both the Va and Via groups have a b.c.c. crystal structure, rhenium has an h.c.p. lattice. None of these metals exhibits a crystallographic phase change upon a change in the temperature. For high-temperature use, a low vapor pressure, low temperature of oxide sublimation, high thermal conductivity and low thermal expansion are important properties. Table 10-2 summarizes some typical data.

599

10.3.2 Chemical Properties Refractory metals are to various extents prone to destruction through gases (Chandler and Walter, 1968; Inouye, 1968), but they have excellent resistance to many chemical agents (Lambert and Droegkamp, 1984; Lambert, 1990). Group Va metals absorb large quantities of hydrogen, and can in consequence be seriously embrittled. Group Via metals remain unaffected. No such effect as hydrogen embrittlement in steel exists. The formation of surface oxides proceeds much slower for chromium than for tantalum, niobium, or tungsten, as shown in Fig. 10-2. Molybdenum and tungsten start to oxidize in air around 300-400°C. Volatile oxides form and evaporate catastrophically above 600 °C in the case of molybdenum, above 800 °C for tungsten, and already above 300 °C for rhenium. Resistance to gaseous hydrocarbons and carbon oxides is defined by the temperatures of carbide formation. Nitrides only form on the surfaces of molybdenum and tungsten under high-pressure nitrogen and ammonia of 1 bar (100 MPa). The damage is limited to brittle layers which form very slowly at the usual temperatures.

0

10 Time (h)

Figure 10-2. The oxidation of refractory metals in air at 1090 °C (Tietz and Wilson, 1965).

600

10 Refractory Metals and Their Alloys

Table 10-2. Important physical properties and mechanical room-temperature properties of pure refractory metals. Group Va

Property V Structure b.c.c. Density (g/cm3) 6.1 Melting point (°C) 1920 + 20 Coeff. of linear thermal expansion a = (d//dr)//(xlO 6 ) 8.3 Specific electrical resistivity at 20 °C (jiQm) 0.19 Approx. interstitial solubility (ug/g), 20 °C for C to H 10 3 -10 4 Start of recrystallization (°C) (1 h anneal) 650 Modulus of elasticity (GPa) 130-150 Poisson's ratio 0.37 Vickers hardness HV10; recrystallized (min.) 70 as-worked (max.) 200 Minimum UTS (MPa) recrystallized 300 DBTT (°C), bending recrystallized (min.) -100

Group Via

Group Vila

Nb

Ta

Cr

Mo

W

Re

b.c.c. 8.6 2470 + 20

b.c.c. 16.6 3000 + 30

b.c.c. 7.2 1920 + 20

b.c.c. 10.2 2620 + 20

b.c.c. 19.3 3380 + 20

h.c.p. 21.0 3180

7.6

6.6

6.2

5.3

4.6

8.4

0.14

0.12

0.12

0.05

0.05

0.17

10 2 -10 4

70-4000

0.1-1

0.1-1

<0.1

<3000

750

1100

700

800

1150

900

100-120 0.40

180-190 0.34

220-280 0.21-0.28

320-360 0.30

390-410 0.30

470 0.29

80 200

80 200

100 300

150 500

250 580

300

300

300

550

350 650 480 °C 600

-200

-269

+ 200

-20

+ 400

Refractory metals are in general resistant to corrosion in various media (salt water, sodium-hydroxide solutions), but those of group Va group are outstanding in their performance. Their stability in dilute, nonoxidizing acids is good, but oxidizing acids corrode the metals severely, especially molybdenum. All of the refractory metals are attacked by sodium and potassium hydroxides, especially when oxidizing substances are present. Their resistance to solids and metal melts (e.g., liquid alkali metals when oxygen is absent) is generally good. Their outstanding corrosion resistance to glass and quartz melts is the reason for their wide application in the glass industry. Of superior resistance against all

1070

glasses are molybdenum-silicon alloys containing a few percent of glass-compatible silicon (Eck, 1984). A medium in which chromium is unsurpassed by conventional materials (such as nickel-based superalloys) in its resistance is vanadium pentoxide, which is present in ashes of all fossil fuels. Refractory metals are not toxic (Goyer, 1990), but their chemical compounds are, to various degrees. In particular, vanadium is reactive enough to convert easily into toxic compounds. The biocompatibility of niobium and tantalum is known to be excellent, which suggests their application as surgical implants (Schider et al., 1985; Zitter and Plenk, 1987).

10.3 Characteristics of Pure Refractory Metals

10.3.3 Recrystallization Behavior

Recrystallization diagrams of refractory metals showing the correlation between grain size, degree of deformation, and applied temperature are similar to those of other metals (see Tietz and Wilson, 1965; Benesovsky, 1979). Table 10-2 includes approximate recrystallization temperatures of pure refractory metals. Figures 10-3 to 10-5 give information on the starting and

0.2

1 Annealing time (h)

10

Figure 10-3. Conditions for complete recrystallization of niobium and tantalum deformed by different amounts (Pionke and Davis, 1979; Lupton, 1991). IOUU

c 1^^

Mo~z±

uoo 1000

ZHM ,MHC

^rjlL

:

•

—

-

TZM

•+1Re

__E__

—

s

Mo

E

Mo-

600 •

i

i

i

i

i

.

•

S i

i 1

i

100

1 10 Annealing time (h)

0.1

Figure 10-4. Start (S) and end (E) of recrystallization of molybdenum and some of its alloys (Eck, 1979).

601

final temperatures of recrystallization, and on their time dependence. The recrystallization of pure metals depends on their actual purity, i.e., on segregations at the grain boundaries. The influence of certain alloying is indicated in the diagrams, and will be further evaluated in the following sections. 10.3.4 Important Mechanical Properties

Characteristic strength and hardness data are quoted in Table 10-2. Figures 10-6 to 10-8 show the mechanical properties of recrystallized, pure refractory metals. The essential feature of these diagrams is the lack of ductility (i.e., zero fracture elongation) of Via metals in a low-temperature range; they break prematurely in a brittle manner so that their strengths, which should be higher than those of Va metals, cannot be exploited to their full potential. This change from ductile to brittle behavior at a certain temperature (the ductilebrittle transition temperature, DBTT) is discussed in Sec. 10.7.6.1. In a medium temperature range between 200 and 800 °C, the strength curves of the Va metals tend to have peaks (see Sec. 10.7.5.2) reaching above the stress levels of the Via metals (these peaks are eliminated from Fig. 10-7 as in the original literature "to avoid confusion"). Slight minima appear in the elon-

— —

-EW-2ThO2

W-26Re UOO

"^^" ———.

W 1000 l

i

l

t

I

I

o U) .—

D

1

I

10

400

s—_. — E~

"^^—s_ 1

0.1

..

I

200

a

— —._ y

~~Ta

o -~" I

n

I

100

Annealing time (h)

Figure 10-5. Start (S) and end (E) of recrystallization of tungsten and some of its alloys (Eck, 1980).

i

i

t

500 1000 1500 2000 Test temperature °C)

2500

Figure 10-6. The temperature dependence of the modulus of elasticity of refractory metals (Tietz and Wilson, 1965; Lambert, 1990).

602

10 Refractory Metals and Their Alloys

gation curves, which correspond to these peaks. The UTS of pure rhenium is superior to that of all the other refractory metals. The fracture-toughness curves in Fig. 10-8 also reflect the transition from ductile to brittle behavior. 1000

Creep properties of pure refractory metals can be demonstrated in stress-rupture plots. Figure 10-9 shows the low creep resistance of unalloyed niobium and tantalum, which is comparable to that of nickelbase superalloys (the diagram also shows the potential for improved creep strength through alloying). Figure 10-10 compares the creep of chromium of different purity levels, and of tungsten tested between 700 and 1200 °C. Creep properties of molybdenum, tungsten, and their alloys are shown in Figs. 10-11 and 10-12 giving the relation between the applied stresses and the testing temperatures that lead to lifetimes of 75 000 hours (8.6 years).

2000

500 1000 Test temperature

1000

Nb-30W-1Zr-0.06C 1200 °C Ni-base MA6000E 980 °C

o

80

Q_

(b)

60 -

Mo *

Nb /

—

•

—

To

-20 -

—^><_ r w

-P

C r / /

•A—^ — ^ To

Nb

J= 100 T3 CD

Nb 980°C ~— Nimonic 105 1000 °C

Q_ <

0

' -200

200 400 600 800 Test temperature (°C)

Figure 10-7. The temperature dependence of (a) UTS, yield stresses, and (b) fracture elongations of recrystallized refractory metals (Tietz and Wilson, 1965; Lambert, 1990). - worked recryst.

/TZM

3000 - '

?FE CD ^

2000 -

Nb

I

o

/

\

j

/ M o j

j /

10 100 Time to rupture (h)

>

1000

Figure 10-9. Creep-rupture curves for niobium, selected alloys, and tantalum (Lupton, 1991). 1000 a

— Cr 99.7 — — Cr 99.97

°c

700

— W 870

°C

W

_t

100 -

- . 700 °C

900 ^C

1200

»

-——^

/J

Q_

'

1

Mo /

r: ^ IOOO •

10

1000

—

Q_

•

—

" ^ - ^ -—

°C

c

900 °C 11f-

-

^ W

\12

05 °C

i

-200

400 0 200 Test temperature (°C) Figure 10-8. The temperature dependence of the fracture toughness for some refractory metals (Marschall and Holden, 1966). The lower branches of the curves represent valid K1C values.

1

°1

0

1

2

3

4

Log [time to rupture (h)]

Figure 10-10. Creep-rupture curves for recrystallized chromium and tungsten tested at various temperatures (Tietz and Wilson, 1965; Eck et al., 1991).

603

10.4 Structure and Properties of Refractory Alloys 1000

50 % of plastic forming. Below this limit, pores still present from the sintered blank deteriorate the ductility. Tungsten is worked below its recrystallization temperature.

0

500 1000 1500 2000 75000-h rupture temperature (°C)

Figure 10-11. Creep-rupture behavior of wrought and stress-relieved molybdenum and its alloys (Hagel et al, 1984). Pure molybdenum recrystallizes in the course of the investigation. 1000 o Q_

0

500 1000 1500 2000 2500 75000-h rupture temperature (°C)

Figure 10-12. Creep-rupture behavior of recrystallized tungsten and its wrought and stress-relieved alloys (Hagel et al., 1984).

Molybdenum and tungsten are rarely used under high-cycle fatigue conditions. Very few data exist and are difficult to compare due to nonconsistent testing conditions (Gold and Harrod, 1979). To strengthen pure metals, only strain hardening and grain refinement can be employed. The heating temperatures during fabrication should be low enough not to induce recrystallization, but in practice, for working temperatures starting at 1400 °C and during intermittant anneals, recrystallization usually occurs in molybdenum. As a consequence, the increase in strength for up to 90% plastic deformation is moderate (Fig. 10-13). A positive effect is the high fracture elongation above

10.4 Structure and Properties of Refractory Alloys Alloying of metals may have different aims: the improvement of physical and chemical properties, or of mechanical properties especially for long-term use at high temperatures. The mechanisms employed to achieve these goals for refractory metals are perhaps more varied than in the case of other metallic materials, although the number of commercially available alloys is low (see Table 10-3). Refractory alloys have not yet been standardized, so "alloy ranges" exist rather than exactly defined alloy compositions. As for all other metals, work hardening is the most basic way to raise the strength. Examples of diagrams which should be similar for all refractory metals possessing

1200 :

1000

0.5

0.7 0.8 1.0 I

1.5

2

2L

o Sheet-rolled A Rod-rolled aRod-extruded • Rod-swaged

.oMo

0.6 0.8 Engineering strain (S o -S)/S o

1.0

Figure 10-13. Strength and fracture elongation of molybdenum and TZM tested at room temperature, as a function of the degree and mode of plastic deformation of the sinter ingot (Eck, 1979).

604

10 Refractory Metals and Their Alloys

Table 10-3. Important commercial alloys and their applications. Alloy

Typical property

Typical application

high strength, good strength/specific weight relation, good corrosion resistance

aircraft and space applications; space applications; nuclear engineering

high strength, good strength/specific weight relation

space applications; nuclear engineering

good corrosion resistance; high strength

chemical process equipment

good resistance against electrical wear, good electrical conductivity

alloys for vacuum contacts

high recrystallization temperature, elongated recrystallized grains, creep resistant, ductile at 20 °C

wires, sheets, and foils for exposure at 1600-1800 °C

low YS-to-UTS ratio resistant against liquid zinc highest ductility, highest strength, electromotive force

wires for lamps and tubes zinc producing industry ductile thermocouples for high temperatures; for welded constructions

high strength, high ductility (MHC and ZHM can be age-hardened)

construction materials for use at 900-1600°C; isothermal forging tools

high recrystallization temperature, elongated recrystallized grains, creep resistant, ductile at 20 °C

construction material for use at 1200-2200 °C

Vanadium alloys V-(5-30)Ti-(5-15)Cr

Niobium alloys Nb-lZr Nb-lOHf-ITi Nb-10W-10Hf-0.14C Nb-10W-2.5Zr Nb-28Ta-10W-lZr Tantalum alloys Ta-(2.5;7.5;10)W Ta-8W-2Hf Ta-10W-2.5Hf-0.01C Chromium alloys Cr-(40-75)Cu Molybdenum alloys Doped (KS-) Mo

Solid solutions Mo-0.05Co Mo-(5-30)W Mo-(5;41)Re

Dispersion alloys TZM Mo-0.5Ti-0.08Zr-0.05C MHCMo-l.2Hf-0.06C ZHM Mo-(0.4-0.7)Zr-(1.2-2)Hf(0.15-0.25)C ODS alloys (Mo-0.7La2O3)

Complex alloys M25WH (Mo-25W-l.2Hf-0.06C) ODS-ZHM-1Y2O3 ODS-ZHM-lLa 2 O 3

highest strength at high temp., brittle construction material for use at at low temp., high strength and creep 1200-2200 °C resistance at high temp.

Tungsten alloys Doped (AKS-)W

high recrystallization temp., elonlamp filaments, metallizing wires gated recrystallized grains, high creep resistance

10.4 Structure and Properties of Refractory Alloys

Table 10-3.

605

(Continued).

Alloy

Typical property

Typical application

Doped W-lThO 2 Doped W-2ThO2

also: low electronic work function

electronic emitters

Solid solutions W-(3-5)Re W-(10;25;26)Re W-lOTi

low-temperature ductility, high-temX-ray targets perature strength, electromotive force thermocouple wires diffusion barrier sputter targets for semiconductors

Dispersion (ODS) alloys W-(l-4)ThO 2 W-0.0ZrO2 W-(2-2)CeO2 W-(l-2)La 2 O 3

low electronic work function, good machinability, high strength the same, but not radioactive as containing alloys

TIG-welding electrodes, plasma generators alternative material to W-ThO2 alloys

Heavy alloys W-7Ni-3Fe W-5.5Ni-2.5Fe W-3.4Ni-l.6Fe W-2Ni-lFe W-7Ni-2.5Cu W-5Ni-2.5Cu W-3Ni-1.5Cu

ductile at room temperature, excellent X- and y-radiation shields, machinability, can be work hardened, counterweights and high-inertia high density (specific weight) materials; artillery penetrators Ni-Cu binder not ferromagnetic

Electrical contact materials W-(10-50)Cu W-20Cu-lNi W-(20-50)Ag

ductile at room temperature, excellent electrical high-voltage contacts, machinability, can be deformed, good welding electrodes electrical conductivity, good resistance against electrical wear

a DBTT, are given in Figs. 10-13 and 10-14. They show a strength increase with plastic work (i.e., "cold work", since the starting temperatures for deforming these alloys are about 1300-1400 °C, and below recrystallization). At first sight it appears unusual that the fracture elongation measured at room and other low temperatures increases with further working simultaneously with the strength. This can be explained to a certain extent by the variation in the DBTT. According to Fig. 10-14, the DBTT of this alloy when deformed by 50% is 300 °C. The material deformed by that amount and tested at 300 °C has therefore become ductile. That the fracture

elongation increases further with progressive cold work seems to be due to the effect of a special thermomechanical treatment (see Sec. 10.7.5.4). For testing temperatures well above the DBTT, the decrease in the fracture elongation with progressing deformation and increasing strength follows expectations. In the diagrams in this chapter, the mechanical properties of the worked materials will often not be directly comparable. Certain alloys are preferentially produced as rod, sheet, or wire, etc., and have therefore experienced different degrees of cold working depending on their prospective application. A shortcoming of some of the

606

10 Refractory Metals and Their Alloys True strain In (S(/ S ) 1.5 2 34 0.5 0.7 1.0 500 DBTT

2000

ouu

80

.^-^UTS /

1 £~]°

3 600 a.

300

.

a)400

U T S ^ " ^ S\

^200 100

20 ~ 800

300°C~

~~~"

UOO

\

n

i

i

UOO °C 15 "

—

i

0 -100

—

—

/

- 20

^-^T^4—^ i

i

500 1000 1500 Test temperature (°C)

0 2000

Figure 10-15. Tensile properties of Nb-28Ta-10W-lZr as a function of test temperature (Gerardi, 1990).

- ^

800 °c — ^

n

--eot.

— stress relieved . -—re crystallized

CD

1000

/

/S

^

10 5

300 °C~-^""'"

n

^

0

20°Cx/

0.2 0.4 0.6 0.8 Engineering strain (S 0 -S)/5 0

1.0

Figure 10-14. Strength and fracture elongation, measured at various temperatures, and the DBTT of W-2ThO2, as a function of the degree of plastic deformation of the sinter ingot (Bildstein and Eck, 1977).

diagrams is also that strain rates are not indicated, but these are mostly expected to be in the range of 10~ 4 -10~ 3 s" 1 . 10.4.1 Substitutional Alloys Vanadium alloys have only recently been developed. Binary alloys with chromium, niobium, and titanium were tried out, but one-phase ternary alloys appear to conform better with the demands for high strength and good corrosion resistance (Harrod and Gold, 1980). The elevated-temperature properties of a high-strength, niobium-base alloy are presented in Fig. 10-15. Common to all such alloys is a ductility minimum between 500 and 1000 °C corresponding to a retarded decrease in the UTS with increasing temperature. Molybdenum alloyed with only 500 jig/g cobalt is the most important material for thin wires in the lamp and electronics in-

dustries. Heat treatment between 1100 and 1200 °C improves, i.e., decreases, the YSto-UTS ratio considerably, so that, for example, helical winding of thin wires for lamp production is facilitated. Mo-W alloys are notable for their outstanding longterm corrosion resistance to liquid zinc (Eck, 1978). Mo-Re and W-Re alloys are renowned for their high-temperature strength and low-temperature ductility, and their superior weldability and corrosion behavior. Rhenium additions also increase the recrystallization temperature (Figs. 10-4 and 10-5). Low rhenium additions improve the properties only marginally, but it is necessary to avoid too high rhenium contents, which cause embrittlement, due to the appearance of the a-phase. Fig. 10-16 shows the hot strength of Mo-(1; 5; 41)Re in comparison to unalloyed molybdenum. Heat treatment changes the mechanical data according to Fig. 10-17. The strength and ductility are superior to other alloys even for heat treatments up to 2200 °C. The elongation of all the alloys has a maximum upon annealing near 1200 °C. The DBTT is always below —70 and above — 180°C in the case of Mo-41Re. Fig. 10-18 shows the UTS and elongation of the worked alloys W-(10; 26)Re. W-26Re in particular is ductile at room temperature, and both exhibit nearly twice the strength of pure

10.4 Structure and Properties of Refractory Alloys

607

1500 1500

:1OOO 00 I— 3

500

en c

500 1000 1500 Test t e m p e r a t u r e (°C)

0

2000

Figure 10-16. Tensile properties of worked Mo-Re sheets of 1 mm thickness as a function of test temperature (Eck, 1985 a).

0

500 1000 1500 Test temperature (°C)

Figure 10-18. Tensile properties as a function of test temperature of W-Re discs forged by 75 %, measured in the radial direction (Eck et al., 1988). 1600 W-26Re

1500

1200 - W-10Re

\

1000 -

Mo-41Re

oo

800

i—

500

^ 5

400 . W-26Re

Mo 1

0

1

1

60 o c

0 1000 2000 3000 1-h heat treatment at temperature (°C)

: "fir

40 -

0

?

o

h-

-40

g

-80

Mo-41Re J ^ I——-27

Figure 10-19. The effect of annealing on the roomtemperature tensile properties of the same materials as shown in Figure 10-18.

]y\

^—•

Mo

Mo-5Re Mo-41Re

T f

>-170

" 1 2 °0 500 1000 1500 2000 2500 1-h heat treatment at temperature (°C)

Figure 10-17. The effect of annealing on the roomtemperature tensile properties and the DBTT of Mo-Re sheets of the same material as shown in Figure 10-16.

tungsten above room temperature. Heat treatments of W-26Re up to almost 3000 °C do not impair the UTS substantially and preserve the good ductility at room temperature (Fig. 10-19). The maximum which develops upon annealing the W-lORe alloy at 1400°C is much more pronounced in Fig. 10-20 for thin wires of the W-5Re alloy.

608

10 Refractory Metals and Their Alloys

-100 ... 0 500 1000 1500 2000 2500 1-h heat treatment at temperature (°C)

Figure 10-20. The effect of heat treatments on the room-temperature tensile properties, and on the DBTT of 0.6-mm wires of tungsten, AKS-doped tungsten, and a tungsten-rhenium alloy (Bildstein and Eck, 1977).

10.4.2 Doped Tungsten and Molybdenum

Pure tungsten wires which were used in early times in incandescent lamps failed by total embrittlement and "sagging" when they recrystallized during service at high temperatures and formed a "bamboo" structure with grain boundaries perpendicular to the wire axis. This failure mode is avoided in ^4^S-doped tungsten which was first produced accidently more than 70 years ago, and has been developed for use above 2500 °C (Pink and Bartha, 1989). The "dopants" are introduced into the oxide powder before it is reduced to metal. A

typical dopant is the so-called "AKS"-dopant, which is a combination of aluminum, potassium, and silicon compounds. After reduction, these elements are incorporated into the tungsten particles. Aluminum and silicon are removed during subsequent washing and evaporated during sintering, leaving only a maximum amount of 50100 jig/g elementary potassium in the sintered rod. Since it is not soluble in tungsten, it ends up enclosed in bubbles (less than 0.1 jam in diameter and visible only after coarsening by annealing above 1900 °C, see Fig. 10-21) which are responsible for the special elongated and overlapping grains which develop upon recrystallization (Fig. 10-22). The mechanism by which the bubbles enforce the formation of the elongated grains, and how the bubbles improve the strength and ductility is described in Sees. 10.7.4, 10.7.5.3, 10.7.5.5, and 10.7.6.2. Since sagging is no longer a problem, this material is also called "nonsag" tungsten. Such wires are drawn to diameters of 5-15 jim, and acquire a UTS of nearly 5000 MPa. For molybdenum, doping with potassium silicates (leaving 150-300 \ig/g potassium, 500-900 |ig/g silicon and 300-500 jig/g oxygen within the metal structure, see Table 10-1) has a similar effect (Eck, 1974, 1979; Fukasawa et al., 1985). Since this KS-doped molybdenum is sintered at lower temperatures than tungsten, a small amount of the additives remains in the form of deformable oxidic potassium-silicon compounds. Recrystallization of doped tungsten starts at 1600-2000°C, that of doped molybdenum at 1600 °C, as is evident from the diagrams giving strength, elongation, and DBTT as a function of the heat treatment (Figs. 10-20 and 10-23). The DBTT decreases with progressive wire drawing according to Fig. 10-24 with a resulting in-

10.4 Structure and Properties of Refractory Alloys

609

sheet (Fig. 10-25). The temperature for complete recrystallization is raised to 1600-2000°C. These alloys are in service above 2000 °C without being destroyed by blister formation. Since they are true dispersion alloys, they will be described in the next section.

Figure 10-21. Scanning electron micrograph of bubbles in AKS-doped rods, deformed by 87 % and annealed 5 min at 1700°C (courtesy of S. Yamazaki).

Figure 10-22. Micro structure of a recrystallized AKSdoped tungsten wire of 180 jam diameter (courtesy of H. Yamamoto).

500 1000 1500 2000 1-h heat treatment at temperature (°C)

Figure 10-23. The effect of heat treatments on the room-temperature tensile properties of molybdenum, TZM, and KS-doped molybdenum wire (Eck, 1974).

crease in the grain-aspect ratio / ( t h e ratio of grain length to width). At very high temperatures (above 1800°C in the case of molybdenum, and above 2800 °C for tungsten), blistering occurs and is caused by silicates or by exaggerated bubble coarsening when the vapor pressure of the enclosed potassium is sufficient to swell up the material. Another group of materials, the only recently introduced ODS alloys of molybdenum, also develop an elongated grain structure upon recrystallization. Compared to AKS- or KS-doped materials, this happens in an Mo-2La 2 O 3 sheet, for example, at a relatively low degree of deformation, as accomplished in a 2 mm thick

100

oo Q

-100

L 2 1 0.5 1 Wire diameter (mm)

10 Grain aspect ratio, f-l/w

100

Figure 10-24. The DBTT of recrystallized KS-doped molybdenum wire: (a) its dependence on the degree of preceding coldwork (Eck, 1980), (b) its dependence on the grain-aspect ratio of recrystallized 1 mm sheets, compared with cold worked (c.w.) molybdenum (Pink and Sedlatschek, 1969).

610

10 Refractory Metals and Their Alloys

Figure 10-25. Microstructure of a 2-mm thick Mo2La 2 O 3 sheet recrystallized at 2300 °C.

10.4.3 Dispersion-Strengthened Alloys One effect of dispersoids, besides direct strengthening, is to retard recrystallization, which is evident from Figs. 10-4 and 10-5. Thus they indirectly improve the high-temperature strength. The most important refractory dispersion alloys are those of molybdenum, although they constitute not more than an estimated 5 % of all the molybdenum alloys. The strengthening agents are carbides of titanium, zirconium, and hafnium, hence the commonly used (but inconsistent) abbreviations TZM, ZHM, and MHC (where "M" stands for the molybdenum matrix). TZM is the most widely used alloy. Carbide dispersion-strengthened molybdenum alloys are often used at moderate temperatures, preferably between 800 and 1100°C. Figure 10-26 shows the results of creep tests performed on TZM at 900 °C during times exceeding three years. Their creep strength is best when they are severely cold worked and subjected to temperatures below that of recrystallization. Recrystallized TZM was reported to behave peculiarly; the curve does not fall monotonously with time, but exhibits a retrogressive tendency at a medium stress.

The response of TZM to aging treatments is too small to be of commercial importance. Newly developed alloys such as MHC and ZHM (for their composition, see Table 10-3) can be aged to achieve maximum strength and ductility. Results from creep tests are listed in Table 10-4. Figure 10-27 shows the excellent strength and elongation of MHC sheets rolled to three different total degrees of deformation with intermittent two-hour strain aging at 1200 °C and heat treatments after every 30 % reduction by rolling (Eck and Tinzl, 1989). A one-hour final heat treatment at 1500°C (see Fig. 10-27) is the standard procedure for obtaining a ductile MHC

TZM arc-cast, 7 5 % worked TZM sintered, 9 2 % worked

£1000

100

TZM arc-cast, recrystallized

n

-

1

1

0

1

i

i

i

1 2 3 4 Log [time to r u p t u r e (h)]

5

Figure 10-26. Creep-rupture curves of TZM tested in helium at 900 °C (Eck, 1979).

1500

0 500 1000 1500 2000 1-h heat treatment at temperature (°C) Figure 10-27. The effect of heat treatments on the room-temperature tensile properties of thermomechanically treated MHC sheet after three different degrees of deformation (Eck and Tinzl, 1989).

611

10.4 Structure and Properties of Refractory Alloys

sheet. The room-temperature strength remains high if practical forming exceeds 75%, which is the minimum deformation for fully dense sheet material. As might be expected, the recrystallization temperature decreases with increasing degree of deformation, lowering the heat-treatment temperature for maximum ductility. In the alloys Mo-(1.8-2.1)Hf-(0.60.7)Zr-(0.23-0.27)C, selective control of Table 10-4. The steady-state creep rate of molybdenum alloys (Eck, 1989; Eck and Tinzl, 1989). Alloy

KS-Mo Mo-lLa 2 O 3 Mo-1Y2O3

Temperature (°Q

Stress

Rate

(MPa)

(h" 1 )

1800 1800 1800

10 10 10

5xl0"5 2.2 x l O " 4 <5.2xlO" 6

1100 1100

450 450

4.2xlO~ 3 2xlO"3

1100 1100

450 450

4xl0~4 l.lxlO"4

1100

450

<10~ 5

Not aged MHC ZHM Aged* MHC ZHM ZHM-lCeO 2 , ZHM-1Y2O3, ZHM-lLa 2 O 3 , ZHM-lThO 2 a

1200 °C for 2h after each of three reductions by 30%.

the maximum strength and high ductility at a chosen service temperature is possible by varying the aging parameters (see Table 10-5) when the Hf-Zr carbide concentrations are high (2.6-3 wt.%). A strength at 950 °C of 750 MPa can be achieved if ZHM is strain-aged at 1200 °C for 2 h after every 30% reduction. Maximum strength at 1450 °C and room-temperature elongations of 15% can be achieved by strainaging for 8 h at 1200 °C after every 50% reduction. Figures 10-11 and 10-12 contain further creep data for molybdenum dispersion alloys, so does Fig. 10-33 in comparison to tungsten, niobium, and tantalum alloys. Data for ZHM and MHC at 1100°C are shown in Table 10-4. ODS molybdenum alloys are doped with the rare-earth oxides La 2 O 3 , CeO2 or Y 2 O 3 (Eck, 1989; Endo et al., 1989; Leichtfried and Wetzel, 1989). They are added before the metal is reduced, or they are introduced to the metal powder before consolidation in amounts of about 0.3-4 wt.%. The addition of these oxides leads to the formation of complex phases consisting, for example, of Y-Hf-Zr oxides when the base alloy is ZHM (Yin et al., 1992). They improve the strength and especially the ductility at room temperature after heat treatment, but may be the reason

Table 10-5. Measures to improve the UTS and the fracture elongation A of ZHM and ODS-ZHM with high carbide concentrations by aging (Eck, 1989; Eck and Tinzl, 1989). Alloy

ZHM ZHM ZHM-1Y2O3 ZHM-1Y2O3

Intermittent (I) and final (F) heat treatment

(I + F) (I + F) (I + F) (I) (F)

1200 °C 1200 °C 1200 °C 2000 °C 1500°C

2h 8h 2h lh lh

1450 °C

950 °C

20 °C A

UTS (MPa)

(%)

1000 1050

2 15

A

UTS (MPa)

(%)

UTS (MPa)

(%)

750 720 820

8 8 5

220 470 350

24 18 19

650

5

530

9

A

612

10 Refractory Metals and Their Alloys

for intergranular failure above 1400 °C. Tables 10-4 and 10-5 give examples of the strength and excellent ductility of molybdenum with ODS additions. ODS molybdenum has become a substitute for K-Sidoped molybdenum. Tungsten can be strengthened by dispersoids using the same principles as applied for molybdenum, but there is little variety in the alloys (see Table 10-3). One type which has been available for some time is thoriated tungsten with l - 4 w t . % ThO 2 . Thoria raises the recrystallization interval to 1300-1500 °C. It improves the strength properties and the ductility (see Fig. 10-14). Since the thoria lowers the electronic work function, this alloy is a favorite material in all those cases where an easy emission of electrons is needed. During use at high temperatures, thoria is reduced by tungsten and the thorium diffuses to the surface forming a film a few atomic layers thick which is responsible for facilitating the electron emission. The disadvantage of thoriated tungsten is its radioactivity. For this reason, ODS alloys of tungsten with rare-earth oxides have become a substitute forW-ThO 2 . A recent development is ODS niobium dispersion-strengthened with TiO 2 and other oxides, which has been suggested for use as human implants (Schider et al., 1985; Gennari et al., 1989). 10.4.4 Complex Alloys A further improvement is achieved by combining substitutional alloying with dispersion hardening. The most common alloys of this type are given in Table 10-3. Group Va metals are strengthened by hafnium carbides, nitrides, and oxides, and their effect is elevated by solution strengthening through tungsten, zirconium and hafnium (Table 10-3).

Some of these complex alloys can be treated thermomechanically so that their strength is raised further (cf. Table 10-4 and 10-5). In the case of M25WH, an addition of 25 wt. % tungsten to MHC increases the UTS to a level (about 750 MPa at 1450 °C) that has not yet been surpassed by other materials, but this alloy is brittle at low temperatures. Optimum combinations of properties for medium- and high-temperature use (cf. Table 10-4 and 10-5) are achieved for dispersion-strengthened molybdenum alloys by additions of rare-earth oxides, e.g., ZHM-1Y 2 O 3 (Eck, 1989). The steadystate creep rate at high temperatures could also be reduced to below 10~ 5 h " 1 for alloys with 1 wt. % CeO 2 , La 2 O 3 , or ThO 2 (Table 10-4). Complex tungsten alloys combine the advantages of rhenium and hafnium carbides (Fig. 10-12). Tungsten alloyed with rhenium, e.g., W-(3-5)Re, can be AKSdoped to develop an elongated grain structure after recrystallization, which opens new fields of application through superior shock resistance. When used for an automobile lamp filament, the doped W-3Rewire reached a life of 3000 h compared to only 1000 h for the undoped (Ogura and Hayashi, 1991). 10.4.5 Refractory Alloys for Electrical Contacts The purpose of combining such divergent metals as, for example, chromium, molybdenum, or tungsten with silver or copper is to create materials for electrical contacts. Thus certain properties are optimized which are incompatible in normal alloys: high electrical and thermal conductivity, high resistivity to chemical reactions, a high melting point for good resistance to burn-off and erosion, good me-

10.4 Structure and Properties of Refractory Alloys

chanical strength, and easy machinability (Goetzel, 1984; Slade, 1986; Shen et al., 1990). These materials have been called "pseudo-alloys" because they are not alloys in the strict meaning of the word. Their components do not react with each other, except perhaps only in a minute layer at the interfaces. Commercial tungsten-copper and tungsten-silver composites range from 10-60 wt.% copper and from 20-50 wt.% silver. Alloying with low amounts of silver or copper (< 15 wt.%) can only be accomplished by P/M. A skeleton of the refractory metal is first produced by pressing and sintering. The interconnecting pores are then infiltrated with liquid silver or copper. It is important that the pores are of optimum size so that the capillary forces are not counterbalanced by frictional forces. Oxide layers may retard the infiltration. Trace impurities in amounts of several hundred }ig/g are also of influence, e.g., nickel improves, while silicon inhibits penetration (Danninger and Lux, 1991). Contacts materials with large or excess amounts of copper (>50 wt.%, i.e., 68vol.%) or silver are differently produced. A mixture of the two powder components is pressed and solid-state sintered at temperatures below the melting point of the low-melting constituent. In such structures the refractory-metal component does not form a contiguous skeleton. Upon heating above the melting point of copper or silver, the alloy would be destroyed in contrast to infiltrated skeleton structures. Such solid-state sintered tungsten-copper can be produced to densities near 92 % of the theoretical. Further densification is possible by plastic forming at elevated temperatures. A typical microstructure of an infiltrated material is shown in Fig. 10-28. Properties of W-Cu/Ag and Cr-Cu alloys are col-

613

Figure 10-28. Microstructure of tungsten infiltrated with 30-wt. % copper.

lated in Table 10-6. Those with silver have superb oxidation resistance. Contacts with copper cost less but should only be used in noncorrosive environments. Contacts with a molybdenum skeleton oxidize more readily and erode faster on arcing. The strength of the composite material can be improved by cold rolling to obtain workhardened sheet material, or by solution strengthening of the soft phase. Before the development of carbon-based materials, another application for such composite metals was as propulsion nozzles and heat shields which are operated close to the melting point of the high-melting constituent (Goetzel and Lavendel, 1965). The low-melting phase can evaporate when subjected to high working temperatures; this cools the high-melting constituent and extends the service life. Porous tungsten with usually 20 vol.% pores and infiltrated with barium aluminate is used for electronic emitters because the electron emission is enhanced through lowering the work function. 10.4.6 Tungsten Heavy Alloys Properties like highly effective attenuation of X- and y-rays, good electrical conductivity, low electrical erosion and welding tendencies, and high specific weight

614

10 Refractory Metals and Their Alloys

Table 10-6. Properties of contact alloys. Property

Contact alloy

Density (g/cm3) Young's modulus, E (GPa) Brinell hardness, HB10 UTSat20°C(MPa) Mean linear coeff. of thermal expansion (wxlO" 6 /K) Specific electrical resistivity (uO) Heat conductivity (W/m K)

W-20Cu

W-35Cu

W-20Ag

W-35Ag

Cu-25Cr

15.4 230 180 400-460

13.5 185 130 340-380

15.6 210 180 240-470

14.0 135 100 300

8.3 110 340

8.5 0.050 134

0.038 155

8.7 0.050 138

0.040 154

13.5 0.040 235

Figure 10-29. Microstructure of a tungsten heavy alloy with 7-wt. % nickel and 3-wt. % iron (courtesy of H. Danninger).

IZUU

UTS

1000 800 -

/ - 30 - 20

0.8 >, '5 0.6 ~ c o o

600

C

- 10 0.A

Ann 90

92 94 96 98 100 Tungsten content (wt.%)

-JO.2

Figure 10-30. Optimum values of room-temperature strength and fracture elongation and their correlation with the contiguity factor (see Sec. 10.7.5.7) for tungsten heavy alloys (Edmonds, 1991).

make tungsten a favorite material for certain applications (see Tables 10-3 and 10-10). In "heavy metals", in which tungsten is the main constituent (Edmonds, 1991), all of these properties can be exploited without having to tackle the inherent brittleness of the tungsten. Heavy metals have full theoretical density right after sintering (see Sec. 10.7.1). They can easily be machined and even deformed. Heavy metals are mainly alloys of 9 0 98 wt.% tungsten and an f.c.c. binder phase of nickel and iron of Ni/Fe ratios of 1:1 4:1 (preferably 7:3, because this composition allows precipitation hardening of the matrix). Nickel improves the mechanical performance, but alloys in which copper is substituted for iron (in Ni/Cu ratios of 3:2-4:1) have better electrical conductivity and are not ferromagnetic. Other alloying metals have been introduced in order to improve the strength and ductility, but they have not gained practical importance. The microstructure of tungsten heavy alloys consists of round tungsten "spheroids" embedded in the binder matrix (see Fig. 10-29). Typical values of strength and fracture elongation are shown in Table 10-7 and Fig. 10-30. They can be seen to vary with composition. The temperature depen-

615

10.4 Structure and Properties of Refractory Alloys

Table 10-7. Properties of tungsten heavy alloys with 90-97.5 wt.% tungsten and a nickel-iron (70:30) binder in the sintered state. Property

Value

Density (g/cm3) 17-18.5 Young's modulus, E (GPa) 350-400 Poisson's ratio 0.28-0.29 Vickers hardness, HV10 270-360 UTS(MPa)at20°C 1000-870 500 °C 650-340 1000°C 260-220 0.2% YS (MPa) 700-600 Fracture elongation (%) 30-10 Compressive strength (MPa) 4500-3500 6300-800 KbtMPa^/mm) Charpy value (notched) 2.8-0.9 Mean linear coefficient of thermal expansion (K" 1 ), 20-800 °C (6.5-5.2) x l O - 6 Specific electrical resistivity (JIQ m) 0.18-0.10

dence of the fracture energy obtained from Charpy tests shows that heavy metals also exhibit a DBTT. In contrast to the sharp transitions known for pure tungsten, those of heavy alloys spread over several hundred degrees, e.g., from - 1 0 0 - + 300°C (German et al., 1984). Phophorus and sulfur act as grain-boundary contaminants, but their effects (see Sec. 10.7.6.3) can be reduced by heat treatments so that fracture elongations of 50% become possible. Heavy alloys can be cold worked to achieve a high UTS at the cost of toughness and fracture elongation (Fig. 10-31). Instead of applying high degrees of deformation, aging after 25% cold working at the low temperature of 500 °C (indicating reactions in the binder phase) was reported, for example, to raise the UTS to 1500 MPa (Penrice, 1984). 10.4.7 Composites Reinforced with Refractory-Metal Fibers

Wires of crystalline or amorphous refractory metals are used to strengthen oth-

er metals, superalloys (Lambert, 1990), or refractory metals (Stephens et al., 1990; Johnson, 1990) in use at high temperatures. When embedded within the matrix, the poor oxidation resistance of some of the refractory metals poses no problem. In comparison to ceramic fibers, their thermal expansion and ductility are closer to those of the matrix materials, which renders them better suited. The strength of all refractory-metal fibers is high in relation to that of most matrix metals at the applied temperatures, which are in general low from the point of view of tungsten fibers. A tungsten-reinforced Nb-lZr alloy was reported to have a ten- and four-fold creep strength at 1400 and 1500 K, respectively, when compared with the plain alloy. Superalloys strengthened with doped tungsten wires of 0.2 mm diameter reach a room-temperature UTS of more than 1000 MPa, and 500 MPa at 1000 °C (Warren, 1989). Fiber-matrix interactions are possible. Chemical reactions lead to the formation of brittle intermetallic phases. Traces of certain metals (e.g., nickel; cf. Sec. 10.7.4), which are deposited on the fiber surface or come from the matrix metal itself, may diffuse into the wire and activate recrystallization of the reinforcement. Below 1000 °C, no such reactions occur. For pro-

uoo

Limit of swaging o/ o 20 %_

1200

Reduction Tungsten content in alloy

,1000

800

0

10 20 Fracture elongation, A {%)

30

Figure 10-31. The variation of strength and fracture elongation of tungsten heavy alloys with composition and degree of deformation (Penrice, 1984).

616

10 Refractory Metals and Their Alloys

tection at higher temperatures, coating the tungsten fibers by chemical or physical vapor deposition, ion plating or ion implantation has been suggested, but these procedures are not very effective. 10.4.8 Refractory-Metal Cermets The name cermet is derived from the words "ceramics" and "metals" and designates a heterogeneous combination of these two types of materials. Cermets based on materials with similar melting points can be produced by pressing and sintering, or by hot-isostatic pressing. The volume fraction of the metallic "binder" can amount to 15-85%. Cermets with high concentrations of refractory metals are materials which are either strong and wear-resistant, or more capable of withstanding chemical attack. In general, they can easily be machined. Prominent examples are the combinations of molybdenum with 15-60 vol.% zirconia. These can be used as sheath material for thermocouples for measuring the temperature of molten steel (Heitzinger, 1974). The strength (about 300 MPa at room temperature) varies little (and even increases) with temperature up to about 1000 °C. The strength can be raised by substitutional strengthening of the molybdenum phase. 10.4.9 Amorphous Refractory Alloys Refractory metals can be quenched into the glassy state by vapor depositing them in the form of thin films onto substrates cooled down to temperatures below 10 K (Johnson, 1990). The pure, amorphous metals crystallize easily, i.e., far below room temperature, but amorphous alloy films can retain their structure upon annealing up to 1000 K. The alloys are of the metal-metal or metal-metalloid

type, e.g., (Nb,Mo) 50 Ni 50 , Mo 3 0 Re 7 0 , Mo 5 2 Ru 3 2 B 1 6 , or W 4 0 Re 4 0 B 2 0 . The elastic moduli of the amorphous metals are in general not too different from those of the crystalline metals. The yield strength and hardness are typically much higher, with a Vickers hardness HV = 2400 for W 4 0 Re 4 0 B 2 0 , for example, and, with a ratio of HV/YS^3 being common for metallic glasses, the yield stress should be 8000 MPa, which places these materials among the strongest known. Since amorphous metals do not have defects like grain boundaries, they are chemically resistant. Sputter deposition of amorphous layers a few micrometers thick onto steel improves its wear and corrosion resistance. Despite their favorable properties, these materials are not yet used commercially.

10.5 Applications of Refractory Metals and Alloys 10.5.1 Statistics The production volume of refractory alloys is small compared to steel or nonferrous metals. Table 10-8 shows the total consumption figures for refractory elements in recent years. Table 10-9 presents a more detailed list giving the percentage used for the different applications. During recent years an average of 8000-90001 of tungsten and 4000-5000 t of molybdenum were produced worldwide per year as pure metal or for alloys, as semi-finished or finished products, which renders them the most utilized refractory elements, unalloyed or as refractory-metal base alloys. An increase in the consumption of tungsten and molybdenum metals and alloys of 2 - 3 % per year is estimated. Chromium is a relatively new metal on the market. While its consumption in various forms is the highest of all the refractory elements

10.5 Applications of Refractory Metals and Alloys

617

Table 10-8. Worldwide consumption of refractory elements. Refractory element (consumption in t/year)

Year

1972 1976 1981 1988 1990

16000

Nb a

Ta

5 200 6000 20000

800

Cr b

2.5 x 106 1000

a

Nb (1989): 250-3001 metal products; contacts).

b

Mo

W

Re

90000 100000 97000 105000

36000 40000 44000

4 5 8.4 15.9

Cr (1990): 75 t metal products (501 sputter targets, 25 t Cu-Cr

Table 10-9. Percentage of refractory elements used for the main products and industries. Product

Element V (1991)

Metal products In steels In superalloys In hard metals For chemicals Miscellaneous

<2 90

Nb (1972)

3 79

Ta (1988)

83 8 6 3

4a

18

Mo (1978)

3-4 83 3 9 1-2

W (1990) Europe

Japan

15

17

Re (1990) U.S.A. 23

10-20

J 625

10-20

1 |20 55 5 5

j 23 58 1 1

5 5

60-70

In titanium alloys.

(see Table 10-8), not more than 75 t were produced in 1990 in the metallic form, and these only for sputter targets (50 t) and Cu-Cr contacts (251). Consumption of vanadium in its metallic form is negligible, and most of it is used in the steel industry for alloying. 250-300 t metallic niobium were used in 1989. Two thirds of the 1000 t of tantalum produced worldwide in 1988 were consumed in the electronics industry, mostly for P/M fabricated electrolytic capacitors. Only about 150-200 t were metal products, primarily wire and sheet. The rhenium consumption in 1988 was over 8 t, estimated to serve primarily for catalytic and alloying purposes. These numbers are expected to grow when new applications

are found and new alloys are developed and substituted for inferior materials. 10.5.2 General Applications

In addition to general remarks in previous sections, Tables 10-3 and 10-10 provide more detailed examples of uses of refractory metals. Special cases will be briefly mentioned here. The chemical resistance of group Va metals makes tantalum in particular a favorite material for the chemical industry. For applications in aviation and for voluminous parts or constructions where their own weight might create additional stresses, strength-to-density considerations

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10 Refractory Metals and Their Alloys

Table 10-10. Application of refractory metals and their alloys. Refractory metal

Application

Vanadium

aircraft industry; corrosion-resistant material for liquid-metal heat exchangers; diffusion barrier for superconductors

Niobium

vacuum-tube getters; rocket skirts, cones and heat shields; missile components; superconductors; chemical process equipment; heat shields in high-temperature vacuum furnaces; nuclear engineering equipment; fusion reactors

Tantalum

electrolytic capacitors; heat exchangers and heaters for the chemical industry; thermometer sheaths; vacuum-tube filaments; chemical process equipment; high-temperature furnace components; crucibles for handling molten metals and alloys; aerospace engine components; surgical implants

Chromium

sputter targets; chromium-copper electrical contacts

Molybdenum

spray metallizing; electrodes and stirring rods for glass manufacturing; filaments and support wires for incandescent lamps; electric-furnace heating elements, boats, heat shields, muffle linings; isothermal-forging tools for superalloys; casting dies; missileand rocket-engine components; high-precision grinding-wheel spindles; pumps, launders, valves, stirrers, thermocouple sheaths for zinc-refining processes; switch electrodes; support and backing for transistors and rectifiers; resistance-welding contacts

Tungsten

incandescent and halogen lamp filaments; anodes and targets for X-ray tubes; semiconductor supports and contacts; electrodes for inert-arc-gas welding; high-capacity cathodes; electrodes for xenon arc lamps, fluorescent and mercury-vapor lamps; automotive ignition systems; rocket nozzles; electronic-tube emitters and anodes; uraniumprocessing crucibles; heating elements and radiation shields; reinforcement in metalmatrix composites; electrical contacts; resistance-welding electrodes; vacuum metallizing; quartz-melting furnaces

Tungsten heavy alloys X-ray and y-radiation shields; aircraft counterweights; self-winding watch counterweights; aerial-camera balancing mechanisms; helicopter rotor-blade balance weights; golf-club weight inserts; artillery penetrators; vibration damping; tools for metal machining Rhenium

filament in carburizing atmospheres; X-ray targets

Mo-35Re

1: W-3.9Re-0.4Hf-0.5C 2: Tct-8W-2.4Hf

500 1000 1500 Test temperature (°C)

2000

Figure 10-32. The UTS-density ratio for several alloys as a function of temperature (Hagel et al., 1984). The curves are directly comparable only at temperatures above where the specific alloys recrystallize because the degree of deformation differs.

may be useful. Figure 10-32 compares tungsten and molybdenum alloys to the strongest tantalum and niobium alloys, and to ODS nickel (Ni-2ThO2) (the data are comparable only above the recrystallization temperature of the specific alloy because the degree of deformation of the various alloys differs). The tungsten alloys W-3.9Re-0.41Hf-0.51C, which is the strongest at high temperatures, and W-HfC are not commercial alloys. Figure 10-33 is a similar diagram showing the temperature dependence of the density-normalized creep stress. Up to service temperatures of

10.5 Applications of Refractory Metals and Alloys 100

10.5.3.2 Porous Metals

Nb-28Ta-10W-1Zr TZM 10

0.1

800

1000 1200 U00 1600 Test temperature (°C)

619

1800

Figure 10-33. Density-normalized stress for 1 % creep in 10 000 h (1.1 years) for several alloys as a function of temperature (Hagel et al, 1984).

1600°C, strain-aged ZHM and MHC are the superior refractory alloys. For the temperature range of 1600 to above 2000 °C, tungsten alloys (e.g., with rhenium) are not surpassed by any other metallic materials, and are important even for the aerospace industry despite their high specific weight. Ceramic- and carbon-based materials are alternatives and should be considered for temperatures above 1600°C.

10.5.3 Requirements for Special Applications 10.5.3.1 The Quest for High Purity

The use of refractory metals in microelectronics depends strongly on reaching purity levels of 99.999 and 99.9999%. Reducing the content of radioactive traces of uranium and thorium to below 5 ng/g, and of lithium, sodium, and potassium to below 0.01 jig/g is possible by P/M means. Sputter targets of W-lOTi, which are used to deposit diffusion barriers in semiconductors, are presently fabricated by P/M to yield a purity level of 99.99%, with impurity traces of molybdenum and oxygen which can be tolerated whilst meeting the demands with respect to thorium and uranium.

The most important use of tantalum is in capacitors for application in electronics (Belz, 1959). Tantalum forms the best anodic films, in contrast to vanadium. It is desirable to press the powders to as low a density as possible, while retaining sufficient green strength for handling. Because the capacitance is directly proportional to the surface area accessible to an electrolyte, a porous pellet structure is desirable. Electron-beam melting purifies the tantalum, thereby allowing capacitor devices to operate at higher voltages. 10.5.3.3 Refractory Alloys for Thermocouples Thermocouples of W-26Re/W-5Re and W-25Re/W-3Re were expected to replace Pt/Pt-Rh couples in the near future (e.g., Peng and Zhao, 1989), but when used for measuring temperatures above 2000 °C, both tungsten-rhenium couples embrittle fast. For long-term use up to 2000 °C and for short-term use above, Mo-41Re/Mo-5Re couples were developed. Before being put into service these thermocouple wires have to be aged to minimize the change in thermoelectric force during high-temperature exposure. The voltage of this type of thermocouple is, like that of the tungsten-rhenium couples, in the range of that of Pt/Pt-lORh. 10.5.3.4 Refractory Metals for Superconductors Unlike the pure metals of groups Va and Via, which are transition-element superconductors, some of their alloys and compounds belong to the group of high-field type II superconducting materials having simultaneously favorable critical transi-

620

10 Refractory Metals and Their Alloys

tion temperatures for superconductivity as well as high upper critical magnetic fields (Warnes, 1990). NbTi alloys have the lowest superconducting characteristics (Kreilick, 1990). Better in this respect are the Al 5 superconductors (Smathers, 1990) of cubic crystal structure of type A3R, with A standing for the metals of groups Va (e.g., V 3 Ga) and Via. In contrast to NbTi superconductors they are, as intermetallic phases, brittle. The third group of superconductors with high upper critical magnetic fields are the ternary molybdenum chalcogenides or "Chevrel" phases (Le Lay, 1990) of the type MMo 6 X 8 , where M is a cation and X one of the chalcogens sulfur, selenium, or tellurium. Amorphous modifications of refractory metals, e.g., Mo 3 0 Re 7 0 or Mo 52 Ru3 2 B 16 , also belong to the group of high-field type II superconductors, but their transition temperatures ( < 9 K) are rather low (Johnson, 1990). For physical reasons it is necessary to surround the superconductor with a metal such as copper, and to keep the diameter of the superconducting wire extremely small (10 jim). Wires containing several hundred or thousand individual superconducting filaments are produced by sophisticated sequences of extruding and drawing. Great demands are imposed on the workability of the materials. The malleable alloy NbTi is easily deformed. The production of A15 superconductors is complicated by the extreme brittleness of the A3B-phase. Extrusion and drawing is therefore done with wires of metallic niobium (or vanadium) inserted in sheaths of CuSn- or CuGabronze. Once the composite has been manufactured, the niobium filaments are transformed into superconducting Nb 3 Sn or Nb 3 Ga (or the vanadium equivalents) with the help of diffusion from the sheath materials during successive annealing. Vanadi-

um is used, in the case of niobium-base superconductors, as a diffusion barrier to prevent the tin migrating into the copper cladding. The Chevrel phases PbMo 6 S 8 ("PMS") and SnMo 6 S 8 ("SMS") are synthesized by powder-metallurgical methods (Grill et al., 1989). Metallic molybdenum, which does not react with the Chevrel phases, is used as sheath or matrix material. In order to optimize the critical current density, it is necessary to pin the flux lines. This is achieved by extensive cold work, which creates grain and subgrain boundaries elongated in the direction of the wire axis, and by heat treatments leading, in the case of NbTi superconductors, to precipitation of the a-titanium phase. 10.5.3.5 Requirements for Use in Thermonuclear Reactors

To achieve maximum efficiency in nuclear power stations, the various elements are operated at high temperatures (Mazey and English, 1984; Smith et al., 1985). Refractory metals are favored because of their strength retention despite neutron irradiation, their low neutron-absorption cross section and short induced radioactivity, and their good corrosion resistance (e.g., to liquid alkali metals) at high temperatures. Irradiation-induced void swelling and elevation of the DBTT are undesirable. Refractory metals have also been considered for application in fusion reactors as "first-wall" materials, i.e., as materials situated closest to the plasma. Besides the necessity to withstand neutron irradiation, there is a further criterion for selection: how they cope with electromagnetic energy. For further details about the behavior of refractory metals under neutron irradiation, see Vol. 10 of this Series.

10.6 Further Processing

10,6 Further Processing 10.6.1 Fastening and Joining 10.6.1.1 Mechanical Fastening

For construction parts such as heating elements, heat shields, and containers used at service temperatures up to 3000 °C, refractory metals are preferably fastened by riveting, or by screws and nuts made of the same material. Riveting is cheaper. At the service temperature the threads of the screws and nuts sinter to the parts they are joining, and cannot be dismounted again. Rivets of molybdenum and tungsten are locally heated by torches to 800-1000 °C before pressing. As a result of this, the rivets are cold worked during assembling and increase in strength, but usually recrystallize during service. 10.6.1.2 Welding

Common to all welding of refractory metals (Lison et al., 1989; Lambert, 1990) is that only "clean" welding methods are applicable. Welding under vacuum or in pure, inert gas prevents embrittling contaminations. Electron-beam (EB) and tungsten inert-gas (TIG) welding are the primary methods. The EB-welded joint is strong since a very small volume is involved which remains fine-grained after rapid solidification. TIG welding enables the joining of large parts. A method which does not create a liquid phase is friction welding. An estimated temperature of 2000 °C is reached near the contacting faces for a few seconds, which is not sufficient for recrystallization. Diffusion bonding is another way of welding without a liquid phase, and it also results in superior joint properties. Diffusion bonding is usually performed while the parts are clamped together, often using hot isostatic pressing equipment.

621

The strength of the joint is the same as that of the unwelded material if the surfaces are clean enough to avoid contamination, and if the bonding temperature is kept below the recrystallization temperature. All refractory metals and alloys can be bonded in this way, even when the parts to be joined are made of different materials. For thin parts, low-energy methods such as laser welding are performed. Unalloyed tungsten and molybdenum tend to have excessive grain growth due to welding, creating welds which are brittle near room temperature. Dispersionstrengthened TZM, ZHM, and MHC can be welded successfully because their carbides retard grain coarsening and strengthen grain boundaries. Large vessels fabricated from TZM have been TIG welded. Creep-rupture limits for a testing temperature of 850 °C for friction-welded, worked TZM do not fall to the lower level of the recrystallized material, and an EB-welded TZM rod of 25 mm diameter tested under the same conditions for periods longer than a year performed better than a recrystallized, unwelded rod (Jakobeit et al., 1985). Discs of TZM were diffusion-bonded under isostatic pressure achieving the same strength as the base material (Eck, 1981). Alloys containing rhenium are the best materials for perfect welding of Via metals. Mo-41Re is recommended if high ductility at and below room temperature is necessary (the DBTT in bending is < -100°C; Eck, 1985 a). After EB welding, it retains its high strength over the full range of temperature after stress-relieving heat treatments. An Mo-41Re rod can be used as the filler material for TIG welding of molybdenum and other alloys. Mo-41Re is expensive, but with any decrease in the rhenium content the strength and ductility deteriorate.

622

10 Refractory Metals and Their Alloys

For welding tungsten-rhenium alloys, W-26Re is the corresponding filler alloy, containing the highest possible amount of rhenium before a-phase formation starts. This alloy is ductile at room temperature, and can be EB, TIG, resistance, plasma, or laser welded. The welding behavior of group Va metals is very different from that of group Via metals. At high temperatures, protection from gases such as oxygen, nitrogen, hydrogen, water vapor, and carbon oxides is necessary. The solubility of these gases is high, and even low contents of interstitial elements have an embrittling effect. Postweld heat treatments are necessary and aging reactions of the weld have to be taken into consideration, especially in the case of alloys. The detrimental effects of impurities become especially serious in substitutional alloys of niobium or tantalum. It is necessary to keep the alloying elements below a certain level so that the DBTT will not be raised to an unacceptably high temperature (Buckman, 1988). 10.6.1.3 Brazing Compared to mechanical fastening and welding, the disadvantage of brazing is the limitation imposed by the melting point of the braze, and in many cases by the temperature of the eutectic formed. Another severe restriction is often corrosion of the braze or the brazed area. Contaminationfree surfaces are a must. The cleanest possible method is vacuum brazing. Brazing gains importance for the fabrication of refractory-metal composites where no other method is feasible; brazing to ceramics, graphite, or semiconductor material (silicon) are accepted joining procedures for high-technology applications. For joints between refractory metals, their alloys and superalloys, and high-

alloyed steels a variety of braze materials are in use (Anonymous, 1980; Lambert, 1990), many of which are commercially available as foils. The choice of braze material depends on later functions of the composite, the partner material, and the brazing equipment. Brazing materials for refractory metals and alloys are designed in such a way that they exhibit good wettability or even solubility in the partner materials without the formation of brittle phases. Upon solidification the grain size should remain fine (this is ensured by eutectic compositions). The temperatures at which braze materials can be used range from 700 °C in the case of silver-copper to 2200 °C for tantalum-hafnium alloys. All braze materials should be free of volatile gases. Nickel is an excellent braze itself due to its good wettability. Nickel, platinum, and rhodium are plated to, or often electrolytically deposited on, the refractorymetal substrate. Brazing materials for high-temperature applications are either refractory metals or alloys based on refractory metals. The brazing of tungsten heavy alloys and contact materials is also a standard procedure. Since the melting point of these composite materials and the service temperatures are low, low-temperature brazing is sufficient. Elements contained in the binder phase such as nickel and copper are the base for brazing alloys. For contact materials, copper- and silver-base alloys are in use. 10.6.2 Coatings Great effort has been made to protect molybdenum, tungsten, and rhenium against the formation of volatile oxides, and tantalum and niobium against taking atmospheric gases into embrittling solution when these metals are exposed to

10.6 Further Processing

elavated temperatures (Anonymous, 1980; Wang and Webster, 1982). Besides protective coatings against gas intake, wear resistant coatings and the formation of layers for brazing are commercially important. Coating with protective layers against oxidation is an old research topic, but the perfect solution has not yet been found. Only some specific problems could be solved. Complications arise from differences in thermal-expansion coefficients between compound constituents of layer and base materials, and from the brittleness of the oxidation-resistant layer. Heating alone may initiate microcracks which grow during any thermic cycling, and the smallest crack causes catastrophic failure at high temperatures. The use of existing coatings is limited to temperatures below about 1400 °C, and they have an insufficient life at pressures below about lOOmbar (lOMPa) in oxidizing atmospheres. They give unreliable protection particularly at edges and corners. A further difficulty is that the temperatures applied for coating may destroy the dislocation substructure of the substrate material and degrade it. 10.6.2.1 Pack and Slurry Cementation Pack cementation based on silicides, and to a lesser extent on aluminides, was first developed for niobium and molybdenum. Depositing and letting a homogeneous slurry of complex silicides react with the refractory metal surface is an improvement over pack siliciding. When properly prepared and applied, these coatings are more reliable even under cyclic stressing. The pack cementation of silicides developed for molybdenum has not yet achieved commercial status. But Si-20Cr-20Fe prepared from fine elemental powders suspended in organic liquids is the standard

623

coating for niobium and its alloys. Methods for applying the slurry are dipping, painting or spraying. After application of about 80 fxm of the slurry, the coating is heated to 1300-1400 °C for reaction bonding and diffusion (Gerardi, 1990). The coating of molybdenum with glass is similar to slurry coating. The glass constituents of the layer are applied as a slurry prepared from compound powders, and are reacted at the glass melting temperature. In the case of smaller parts, they are coated by direct reaction in a glass container. The glass composition and the details of procedures are mostly confidential. 10.6.2.2 Electrolytic Coating For protection against oxidation and as a coating for further welding or brazing processes, platinum layers are deposited onto molybdenum using melt electrolytic processes. After diffusion annealing, adherence of the layer is excellent and the coated molybdenum rods can be further processed to wire. The layer thickness is limited by the process itself and by the price of platinum, and usually does not exceed some tenths of a millimeter. Molybdenum electro-plated with platinum or platinum-rhodium is widely used in the glass industry for operation in oxidizing atmospheres. Diffusion of the layer element into the base refractory metal is the main obstacle for successful long-term performance at high temperatures. More widely used for molybdenum than for tungsten are aqueous electrolytic processes to form layers of nickel, silver, gold, rhodium, tin, and copper. Very often a nickel flash improves the adherence of subsequent layers. These coatings are necessary for brazing operations.

624

10 Refractory Metals and Their Alloys

10.6.2.3 Flame and Plasma Spraying

One of the largest areas of consumption of molybdenum is for thermal spray coatings (Clare and Crawmer, 1982). Flame spraying is used for wear-resistant coatings of steel, primarily for piston rings. Molybdenum wire is sprayed through an acetylene flame, melting the molybdenum and thrusting partially oxidized droplets onto the substrate. The resulting wear-resistant layer contains pores which in service act as reservoirs for the lubricating oil. Molybdenum oxides which develop during the process strengthen the layer to a hardness of HV10>1000. Plasma spraying in air (APS) of molybdenum starts with powder granules. Due to the high plasma temperatures, tungsten can be plasma-sprayed equally well, but in contrast to molybdenum, the plasma spraying of tungsten is of limited importance. Iron- and/or nickel-based alloy powders, or elemental blends containing 20 wt. % chromium and 5 wt. % aluminum without or with an addition of up to 1 % yttria [the so-called M-CrAl(Y), where M stands for iron or nickel] are effective as oxidation protection if plasma-sprayed to molybdenum. Glass-melting electrodes can be successfully protected against burner gases during the start of the furnace operation, which lasts for days, until the glass is molten. An improvement to APS, which oxidizes the layers, is plasma spraying in vacuum (VPS). The oxide content of the coatings can be minimized and the densities of the layers increased. Coatings with densities between 95 and 98% of the theoretical and produced by VPS are standard, but higher densities are difficult to achieve. The best oxidation-resistant VPS coatings for molybdenum are again based on M-CrAl(Y).

10.6.2.4 Chemical and Physical Vapor Deposition

Chemical vapor deposition (CVD) and later physical vapor deposition (PVD) processes have stimulated the coating technology for refractory metals. Oxidation-resistant coatings 10-20 \xm thick can be deposited in dense layers when the coating compositions are based on silicides of refractory metals such as MoSi 2 , WSi 2 , CrSi 2 , NbSi 2 , and TaSi2, or SiC. Compositions such as Si-20Cr-20Fe, known from slurry coating, are deposited more easily by PVD processes than by CVD. At present, the layer thickness obtainable from slurry coating makes this method superior to PVD coating. New research on boroncarbide coatings which are applied to refractory metals by means of CVD techniques appear promising (Linke et al., 1992). Platinum and iridium layers are deposited by PVD processes such as magnetron sputtering and ion plating. Deposition of rhenium layers is preferably performed by PVD. These coatings serve as intermediate layers against carbide formation, for instance for refractory-metal-graphite composites. CVD of tungsten and tungsten-rhenium alloys has reached commercial importance for the production of X-ray targets. Tungsten alloyed with rhenium is CVD deposited onto graphite discs of diameters between 100 and 150 mm. For this application final layers of 1 mm thickness are necessary. The CVD process is also used for the production of tungsten and lowalloyed tungsten-rhenium crucibles and other shapes. Decomposing W-Re fluorides or chlorides has to be closely controlled to achieve a fine isotropic grain structure. Densities over 99 % can be obtained. CVD-deposited tungsten is now

10.7 Fundamental Mechanisms

produced to such perfection that general - not CVD-specific - properties are investigated using it (Subhash et al., 1994).

10.7 Fundamental Mechanisms 10.7.1 Activated and Liquid-Phase Sintering

Chapter 4 by Arunachalam and Sundaresan in Vol. 15 of this Series provides general information on sintering. While activated sintering of refractory metals is "micro"-alloying with the object of facilitating sintering, liquid-phase sintering is a way to simultaneously conduct sintering and alloying (German, 1985). The temperatures necessary for comparable amounts of densification can be drastically lowered (from 1800 or 2500 to 800 °C in the case of tungsten) by "activated sintering", which is achieved by small additions (0.2-2 wt.%) of the elements cobalt, platinum, nickel, and palladium (in order of increasing effectiveness), and others which accelerate the atomic-transport mechanism (Li and German, 1984) when they are segregated on the surfaces where two particles are joined (Gessinger and Fischmeister, 1972). The possible formation of brittle phases and contamination of the base metal are severe disadvantages. Activated sintering is not used to produce commercial products because high-temperature furnaces are available. Most of the guidelines for effective activated sintering are equally valid for liquidphase sintering, which is used to produce, for example, tungsten heavy alloys. During the liquid-phase sintering period, the nickel-iron binder is enriched with tungsten (taken from the tungsten particles of the powder mixture) in accordance with its good solubility at the appropriate sinter temperature. During cooling, when the

625

solubility is diminishing, tungsten will be re-precipitated at tungsten-liquid interfaces, and the tungsten particles will in due course assume the shape of a "spheroid", which is mostly a single crystal, containing small concentrations of the other alloying constituents (Fig. 10-29). 10.7.2 Interstitial Solubilities

The solubilities of interstitially dissolved elements have important consequences. Excess concentrations which are not dissolved cause differences in the mechanical properties of group Va and Via metals. Several atomic percent of interstitials are dissolved at room temperature in the Va metals, and only less than 0.1 or 1 jig/g in the Via metals (Table 10-2, Fromm and Gebhardt, 1976). Within one group, the metal with the lowest atomic number dissolves most. The amounts which can be dissolved in the crystal lattices of chromium, molybdenum and tungsten are smaller than the concentrations usually found in commercial products, while vanadium, niobium, and tantalum are capable of dissolving much more than is commonly present (Table 10-1). In mixtures of group Va and Via metals, which are fully soluble in each other, the interstitial solubility is not notably affected. In alloys with the "reactive" group IVa metals the solubility of the interstitials is possibly increased, and the excess concentration is bound in the form of second phases. Of particular interest is the interaction of rhenium with the interstitials in the substitutional group Via alloys. The ternary phase distribution shows that such alloying increases the interstitial solubility (see Fig. 10-34c; Savitsky et al., 1977). Internal-friction experiments (Booth et al., 1965) and field-ion microscopy (Novick

626

10 Refractory Metals and Their Alloys

2 at.% Zr

ture. The dispersoid must have good thermal stability at the temperatures of application. This is ensured to a certain extent when their melting point is high (TiC, ZrC, and HfC have melting points above 3000 °C). Any high-melting particles can in principle be used, although their tendency to coarsen during exposure makes them better or less well suited. Below 1800°C, oxides coagulate with greatest difficulty, and nitrides are more stable than carbides (Savitsky et al., 1982).

3

10.7.3.1 Carbide Particles 12

at.% Ta

20

60

at.% Re Figure 10-34. Phase distribution at 2000 °C for the ternary systems (a) W-Zr-C, (b) W-Ta-C, and (c) W-Re-C (Savitsky et al, 1982).

and Machlin, 1968) demonstrated similarly that the solubility of oxygen is improved by Re-O clustering within the lattice. Interstitials are thus scavenged from dislocations and grain boundaries, and redistributed within the lattice. 10.7.3 Introducing Dispersed Second Phases The best distribution of dispersed particles can be achieved through alloying with elements that are dissolved at high temperatures, and which precipitate due to a solubility that decreases with falling tempera-

To understand precipitation, it is necessary to know the ternary phase diagrams for metals of group Via with the most promising alloying elements of groups IVa and Va, and with rhenium of the group Vila. Representative of these groups are the systems W-Zr-C, W-Ta-C, and W-Re-C, as shown in Fig. 10-34. Characteristic differences exist in the ternary phase diagrams with respect to the alloying metal concentrations necessary to stabilize the carbides MeC: 0.6 at.% zirconium or 0.8 at.% hafnium versus 4 - 6 at.% niobium or 7-8 at.% tantalum. The effect of alloying in molybdenum alloys is similar (Trefilov and Moiseev, 1978). Other phases in equilibrium with the a solid solution (in the tungsten-rich corner) at the given temperatures are W 2 C and the ternary carbides with variable composition W 4 Re 3 C and W 3 Re 4 C (it-phase), and the a-phase in the W-Re system. WC is only stable at low temperatures, and cannot be obtained as a dispersoid. W 2 C is not a highly effective dispersoid because it tends to coarsen more readily than MeC. The conditions for the formation of carbides (as estimated from annealed alloys containing various alloying metals) are documented in Fig. 10-35. HfC is the most

10.7 Fundamental Mechanisms

2200

Maximum solution temperature

Abundant medium rare abundant medium rare

Relative amount Figure 10-35. Range of existence of carbides in dispersion alloys of molybdenum (Ryan and Martin, 1969 b).

stable. Besides the carbides of the alloying metal, MeC, Mo 2 C appears at higher temperatures in a limited range. Increasing the ratio of the alloy-metal-to-carbon concentration, Me/C, e.g., by adding the alloying metal, raises the stability range of MeC and restricts the range of the Mo 2 C. This effect continues until a concentration is reached which suppresses the formation of Mo 2 C altogether (Chang, 1964; Ryan and Martin, 1969 b). The border ratio appears to be 3 through 5 for Me being titanium, 2.5 for zirconium and 2 for hafnium. Alloys with such an Me/C border ratio have the better mechanical properties because they are void of Mo 2 C precipitates which, like W 2 C, tend to coarsen, and are therefore not good hardening agents. There also threshold ratios of Me/C, below which no carbide can be formed. The limiting ratios for ZrC and HfC are a little above the stoichiometric ratio of 1. TiC requires a high ratio of 2.5, and ratios > 5 are necessary to form NbC or TaC in molybdenum. In contrast to such experimental experience, theoretical considerations advocate neither too low nor too high Me/C ratios, but additions complying with the correct stoichiometric ratio of the carbide MeC in order to produce alloys with optimum properties (Wadsworth, 1983). There is

627

some experimental justification for such demands from certain properties, but equally there are results from the same material supporting the previously quoted demand Me/C > 1. Increasing both Me and C while keeping their ratio constant does not necessarily increase the amount of effective precipitates. At the annealing temperature only a certain amount of both elements is dissolved and can be utilized for optimum precipitation. The excess MeC, which remains undissolved, may coarsen during the solution treatment and impair the mechanical properties. 10.7.3.2 Other Particles Alloying through nitrides obeys the same rules as introducing carbides (Ryan and Martin, 1969a,b). Although the results are promising, it has not attracted more than scientific interest. The use of oxides as the strengthening agent was described in Sec. 10.4.3. High-melting oxides are considered to be incapable of age hardening, but combinations with age-hardening carbides are possible. Both internal nitriding and internal oxidation have been tried (Wilcox, 1968; Gaal et al., 1989). The conditions under which submicroscopic bubbles of 0.1 jim diameter and less develop in doped tungsten and act as dispersed "particles" are described in Sec. 10.4.2 and in the following. 10.7.4 Recrystallization in Alloys Dispersion alloys increase the strength in a twofold way: they stabilize the asworked dislocation structure so that recrystallization occurs at higher temperatures. The strength gain from the strainhardened structure is greater than that from the interaction between the dislocations and the dispersoids.

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10 Refractory Metals and Their Alloys

Doped tungsten exhibits a most unusual and anomalous recrystallization behavior, unknown to this degree in any other metal. The cause of this is the potassium-containing bubbles which originate from the doping additions (Horacsek, 1989 a; cf. Sec. 10.4.2). After heavy mechanical working and heat treatments they are aligned like strings parallel to the wire axis (Fig. 10-21). They restrict the movement of grain boundaries in the direction perpendicular to the wire axis, so that the recrystallized grains are elongated and interlocking (Snow, 1989; Fig. 10-22). This doping shifts the recrystallization temperature from the common 1000-1200°C for unalloyed tungsten to 2000 °C and above. Since in fact any dispersoid in any metal is capable of creating a more or less elongated recrystallized grain, the ODS alloys of molybdenum, and to a much lesser extent W-ThO 2 , also acquire elongated and interwoven grains after recrystallization. One type of recrystallization, with disastrous consequences for doped tungsten, is impurity- or chemically-induced recrystallization (Gaal, 1989). Surface contamination by the same elements which activate sintering cause doped tungsten wires to recrystallize at the same low temperatures as wires of undoped tungsten. With activator elements facilitating the diffusion, the bubble structure is apparently not as effective at pinning the grain boundaries as in pure potassium-doped tungsten. Another reason may be that the dissolution and/or coalescence of the bubbles becomes easier. 10.7.5 Micromechanisms of Deformation and Strengthening

Metals deform in a distinct temperature and stress range by a series of mechanisms of which one is always dominant (see Vol. 6 of this Series). For b.c.c. refractory

metals the relevant mechanisms are: (a) the double-kink mechanism (which is responsible for the steep increase in the strength values with decreasing temperature below 0.15 Tm), (b) dislocation climb (either low-temperature or high-temperature dislocation creep) at lower applied stresses, and (c), above 0.5 7^ and at very low stresses, deformation controlled by diffusional (i.e., Newtonian viscous) material transport (either Coble or NabarroHerring creep). The distribution of the mechanisms for refractory metals was demonstrated by Frost and Ashby (1982) in deformation-mechanism maps. Such maps give an approximate idea of the material behavior under conditions of various stresses, strain rates, and temperatures, but they are modified by dynamic recovery and recrystallization. Additional deformation mechanisms in pure refractory metals and alloys (reviewed by Wilcox, 1968) can be sources of further deviations. The slip characteristics of single crystals of refractory metals at the temperatures of the plateau region - the three-stage hardening - are very similar to those of f.c.c. metals. However, in the low-temperature region, b.c.c. metals exhibit a unique, pronounced orientation dependence of yielding and an asymmetry of slip in tension and compression; also, large deviations from Schmid's law of the critical resolved shear stress are observed (Kubin, 1976; Sestak and Seeger, 1978; Christian, 1983). There are certain differences between single crystals of Va and Via metals. In group Va metals, at temperatures below 150 K, the dislocations glide in "anomalous" slip planes which are still less favorably oriented. This behavior is sensitive to the impurity content, and is not accompanied by an asymmetry effect. Slip in molybdenum single crystals is, on the other hand, affected by surface reactions. For

10.7 Fundamental Mechanisms

polycrystalline technical metals, these single-crystal properties are not easily recognized as relevant. In polycrystals of h.c.p. rhenium the temperature dependence of the lattice resistance below 0.15 Tm is similarly steep as that of b.c.c. metals, but it is obscured by the much greater tendency to twinning. Strain hardening is severe even after low degrees of deformation. This is also due to twinning, since twin boundaries act as obstacles for the dislocation motion. On the whole, however, twinning can provide good deformability. Under conditions of compression (rolling, extrusion), in contrast to tension, the basal planes (which are the slip planes) rotate to positions parallel to the rolling plane, and a compressive force acting on them then initiates mechanical twinning (Carlen and Bryskin, 1992). Amorphous alloys deform homogeneously by Newtonian flow when the temperatures are low and the applied stresses are smaller than YS/100. They deform inhomogeneously by localized shear bands when the temperatures are high in comparison to the glass-transition temperature, at stresses higher than YS/50 (Johnson, 1990).

629

phases which improve the strength especially in the high-temperature region (Klein and Metcalfe, 1973). The binder phase of tungsten heavy alloys, which contains some dissolved tungsten or possibly other metals, is also solidsolution strengthened. Substitutional alloying may instigate solution softening in certain metals (see Sec. 10.7.6.1). In ternary systems a synergistic strengthening is observed. 10.7.5.2 Dynamic Strain Aging Additional strength contributions, often in the shape of humps in strength-temperature diagrams, arise from the interaction of dislocations with dissolved interstitial and substitutional impurity and alloying atoms at temperatures between 0.15 7^ and 0.4 or 0.5 Tm. This is very important because at these temperatures the strength of the pure metal is already low. The ductility is impaired as a consequence of this dynamic strain aging. Dynamic strain aging is manifested in many refractory metals and alloys (Tietz and Wilson, 1965; Pink et al., 1985; Pink, 1986). An example of dynamic strain aging around 0.2 Tm is shown in Fig. 10-36 for

10.7.5.1 Solid-Solution Strengthening Binary alloys of niobium, tantalum, molybdenum, and tungsten form continuous series of solid solutions with each other. Only those built of alloy elements with an approximately 5 % difference in atomic size exhibit a high degree of solid-solution hardening (Tietz and Wilson, 1965; Predmore et al., 1972; Klopp, 1975). Rhenium is an important solid-solution strengthener. The improvement due to titanium, zirconium, and hafnium is only partly based on their role as solid solutions. These elements have a tendency to form second

400 o Q_

200

Ta, ~100|jg/g interstitiais,

0

400 800 1200 Test temperature (°C)

1600

Figure 10-36. Dynamic strain aging caused by interstitial elements in recrystallized tantalum (Pink, 1968), and by substitutional elements in a recrystallized molybdenum alloy (Chang, 1964).

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10 Refractory Metals and Their Alloys

commercial tantalum containing a total of about 100 fig/g of interstitial impurities. No such humps appear for molybdenum and tungsten. This supports the diagnosis that the effect is caused by dissolved interstitials; low-temperature dynamic strain aging is prominent in metals of group Va due to their high solubilities for interstitials, and it is in general missing in metals of group Via with their extremely low interstitial contents (the effect appears in chromium with its higher solubility for interstitials). Dynamic strain aging in the higher temperature region from about (0.3-0.5) Tm is caused by substitutionalinterstitial complexes or by substitutionally dissolved alloying elements. The effect also appears in alloys which are mechanically alloyed with oxides when they dissociate during heat treatment and release their constituents into solution. A special manifestation of dynamic strain aging is the Portevin-Le Chatelier effect. In the temperature range where the peak in the strength-temperature curve has a positive slope, deformation proceeds inhomogeneously. There is ample evidence for such inhomogenous flow due to interstitials in metals of group Va and due to substitutional atoms in both groups Va and Via alloys. The role which dynamic strain aging may play in thermomechanical treatments is addressed in Sec. 10.7.5.4. 10.7.5.3 The Effect of Dispersed Second Phases Adding nonmetallic second phases to the metal powder by mechanical mixing before sintering can lead to some strengthening. Stable particles are not remelted during consolidation and sintering, and should preferably be up to 2 jjm in size. They are circumvented by a by-pass mechanism. If larger than 5 jim, they are rather

a cause for weakening. Unstable particles may react with the matrix during processing, which can lead to age hardening and to dynamic strain aging. A certain amount of strengthening can be achieved in eutectica or binary systems when the composition is close to the solubility limit due to the formation of secondary precipitation (e.g., cr-phases in the W-Re system). However, instead of a beneficial effect on the strength, the consequence is often a serious deterioration in the ductility because metal compounds in binary systems tend to coagulate. Age hardening is the most effective strengthening method for conditions which favor deformation by dislocation glide. In most of the high-strength refractory alloys, small coherent particles (1-100 nm) must be cut. Bubbles in non-sag materials also act as dispersed obstacles against dislocation motion (Pink and Gaal, 1989), causing the high-temperature strength to be higher than that of undoped metals. Overcoming the dispersed phases within the matrix does not cancel the basic lowtemperature (i.e., double-kink) mechanism. Its theoretical and the experimentally determined parameters of thermal activation agree over a wide stress range. However, in the case of group Via metals and alloys with and without second-phase additions the enthalpies of deformation in the lowstress range, and thus the extrapolated values of the total activation enthalpies at zero stress, are significantly larger (Dorn and Rajnak, 1964; Pink, 1986), which, it was suggested, stems from a nonconservative movement of jogs created at inclusions. That commercial molybdenum or tungsten behave in the same way as their alloys indicates that they must also be considered as "dispersion-strengthened", which may be justified considering the extremely low solubilities for interstitials.

10.7 Fundamental Mechanisms

In high-temperature deformation of refractory metals, the presence of dispersed obstacles also leads to irregularities in the magnitudes of the activation enthalpies and stress exponents n of the creep equation (Pink, 1986; Pink and Gaal, 1989). The enthalpies, which should indicate core or volume diffusion, are much higher when no care is taken to reduce the applied stress in the creep equation by a back stress (see Vol. 6 of this Series). Equally, the stress exponent can increase when the correction term is not taken into account. A prominent example is bubble-strengthened, nonsag tungsten with n values as high as 50. 10.7.5.4 Thermomechanical Treatments

The term "thermomechanical treatment" comprises all mechanical operations to deform a metal plastically in connection with the application of heat (Perkins, 1968; Trefilov and Milman, 1989). Such operations are, e.g., "cold work", when deformation occurs at lower temperatures, and "warm" and "hot" work below and above the recrystallization temperature. Each deformation mode creates a characteristic dislocation structure of specific density and distribution. Cold or warm work increase the strength and decrease the DBTT. When the work hardening has reached a certain degree, further deformation may become difficult, and intermittent annealing will become necessary. While the curves of the strength data fall and those of the DBTT increase with increasing annealing temperature in the recrystallization temperature range, the fracture elongation passes through a maximum (see, e.g., Figs. 10-17, 10-20, and 10-27). At the start of recrystallization the glide distance within a grain grows, and the fracture elongation becomes larger. But with progressing recrystallization, due

631

to a further increase in the annealing temperature, the grain grows, accompanied by a loss in ductility (i.e., an increase in the DBTT), leading to a sharp drop in the fracture elongation. This ductility maximum may be considered as a case of "transformation plasticity". The stability of the deformation substructure is important for the strength at elevated temperatures; however, it is destroyed by recrystallization. A certain amount of stability is guaranteed when dislocations are decorated with dispersed particles. In principle this is just a description of the impediment to recovery and recrystallization in dispersion alloys. However, the segregation of precipitates at dislocations can also be induced by intentional aging after mechanical working. Such "strain aging", carried out as an intermittent annealing between steps of mechanical working, is a thermomechanical treatment in a more specific sense. Examples of strength gains in alloys have been quoted in previous sections (cf. Table 10-4 and 10-5). The procedures employed are based on experience and have not yet been scientifically explored. It is possible that the outcome is augmented by dynamic strain aging, which occurs during working sequences at certain temperatures (Eck and Pink, 1992, 1993). The result should be an improved cell structure as well as a dislocation structure stabilized by precipitates that may originate in the course of the inhomogeneous deformation during dynamic strain aging. 10.7.5.5 Grain-Size and Grain-Shape Strengthening

For refractory metals, too, grain size has the expected influence (Begley et al., 1968; Eck, 1985 b). At low temperatures, where the Hall-Petch law applies, a fine grain is

632

10 Refractory Metals and Their Alloys

desirable. However, at high temperatures (about 0.7 7^) and low stresses, where deformation is controlled by diffusional transport of matter, the behavior is reversed so that a coarse grain retards creep (see Vol. 6 of this Series). In metals with a recrystallized non-sag structure a further reduction in the hightemperature creep rate (see Fig. 10-37) is due to a grain-shape effect which influences the diffusional flow. Because the diffusional accommodation is complicated in materials with interwoven grains, the grain-aspect r a t i o / ( = grain length/width) has to be included in the high-temperature rate equation (see Pink and Gaal, 1989); the strain rate is inversely proportional to this/ 2 (where z is a constant depending on the type of diffusional flow). This grainshape strengthening is of practical importance at extremely low applied stresses, for example, when tungsten wires used at high temperatures in light bulbs are deformed by their own weight. A deformation-mechanism map (Fig. 10-38) demonstrates how, to maintain a given creep rate, an increase in/allows the lamp filament to be exposed to much higher temperatures at a given stress level. 10.7.5.6 Rhenium-Alloyed Refractory Metals: Deformation Twinning

In refractory as in other b.c.c. metals, twinning is encountered during deformation at low temperatures. Alloying of group Via metals with rhenium or, generally, with metals of groups Vila and VIII promotes excessive twinning (Booth et al., 1965). It occurs with less frequency at high temperatures, yet up to 800 °C, and it was considered to be one reason for the good workability of these alloys. In the low-temperature range, the apparent "yield stress" is independent of temperature down to 77

0 10 20 Grain aspect ratio, f=(grain length/width)

Figure 10-37. The dependence of the steady-state creep rate on the grain-aspect ratio of recrystallized KS-molybdenum (Fukasawa et al., 1985).

-3.5 -

L.T D.C.

o-Vs^

H.T.D.C.

f1

r~

C.

» \

10 zO

\

-4.5 " Coble Creep \1.5

-5.0 1000

i

t

i

\5

\iO

i

2000 Temperature (°C)

3000

Figure 10-38. Deformation-mechanism map demonstrating the effect of a non-sag structure on the creep rate (Pink and Gaal, 1989). L.T.D.C. is low-temperature dislocation creep, H.T.D.C. is high-temperature dislocation creep; ft is the shear modulus.

K. "Yielding" in this case is not a normal plastic deformation process, but it is caused by a series of twin formations accompanied by stress drops. The yield stress is the stress required to form these twins. Higher degrees of straining are accomplished by true plastic deformation, and the UTS exhibits the normal temperature dependence of b.c.c. metals. 10.7.5.7 The Deformation of PseudoAlloys and Tungsten Heavy Alloys

Pseudo-alloys are deformed to gain strength and resistance against wear when utilized as contact materials. The strength of the tungsten can be fully exploited down

10.7 Fundamental Mechanisms

to low temperatures when a tungsten skeleton is infiltrated and surrounded by a ductile binder. Deformation in heavy metals depends on the same principle (German et al., 1984; Edmonds, 1991). Each tungsten spheroid should be completely surrounded by the binder phase, i.e., the structure should be characterized by a low contiguity factor (the ratio of the number of boundaries between two tungsten spheroids to the number of tungsten/matrix boundaries). This factor depends on the volume fraction and on the wetting behavior between the binder and the tungsten. It is necessary to build up a high hydrostatic pressure around the tungsten spheroids in order to avoid premature brittle fracture through the tungsten grains. The strength and work-hardening behavior of the binder in relation to the strength of the tungsten spheroids is also of influence. The uncontrolled variation of all these factors may be the reason why the results of deformation experiments using tungsten heavy alloys have often been ambiguous, testifying to an exclusive initial contribution from either the tungsten grains (Krock and Shepard, 1963; Ekbom, 1988) or the binder phase (Ekbom et al., 1988). 10.7.6 Micromechanisms of Fracture Gandhi and Ashby (1979) plotted the available information for the refractory metals niobium, tantalum, chromium, molybdenum, tungsten, and rhenium on fracture-mechanism maps displaying the distribution of fracture modes as they vary with testing conditions (see Chap. 12 by Riedel in Vol. 6 of this Series). They all look basically the same.

633

10.7.6.1 The Ductile-Brittle Transition Temperature (DBTT) The field of brittle fracture preceded by microplasticity in fracture-mechanism maps is of great practical importance. In b.c.c. metals where the strength increases with decreasing temperature below 0.15 Tm, the yield stress will eventually equal the fracture strength. While fracture above this temperature is ductile, fracture below it will be brittle (the temperature where the behavior changes is the DBTT). In group Va metals brittle fracture occurs by cleavage, i.e., it is mainly brittle transcrystalline. In contrast, metals of group Via with a pronounced grain-boundary weakness will fracture mainly in an intercrystalline mode, because the increasing yield stress first reaches the level of the grainboundary strength. Special brittle-fracture modes characteristic of heavily-worked molybdenum or tungsten sheet or wire are "delamination" or "splitting" (Perkins, 1968; Gaal, 1989), and "45° fractures". Delamination occurs in planes parallel to the sheet surface, and splitting in the longitudinal direction of wires in punching, shearing, or bending operations. Cracks inclined at 45° to the rolling direction and perpendicular to the sheet surface destroy a sheet at the slightest impact. In all these cases the fractures are related to the pronounced texture in such materials, i.e., to the orientation of contaminated cleavage planes ({100} planes in b.c.c. metals). For 45° fractures a (111) component must exist, in addition to a dominant (001)[110] texture. This recrystallizes earlier than the (001) component, and the cracks are initiated at the contaminated boundaries of the new grains but continue to grow as demanded by the main texture (Neges, 1994).

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10 Refractory Metals and Their Alloys

"Ductility" is a term which assumes different meanings depending on the group of materials it describes. It generally denotes the elongation; "high ductility" is synonymous with a large fracture elongation. When talking about refractory metals, ductility is often understood as the condition created by the appearance of the DBTT; a metal has a "good ductility" when its DBTT is situated at a low temperature. Influencing the DBTT Auger spectroscopy carried out on grain and fracture surfaces of group Via metals has established the role of grain-boundary contaminants on embrittlement. Nitrogen segregations in chromium, and oxygen, and, to a lesser extent, carbon segregations in molybdenum, are especially dangerous. Several procedures can influence positively the relation between yield stress and grainboundary strength in group Via metals (Tietz and Wilson, 1965; Perkins, 1968; Stoloff, 1969). Grain boundaries are stronger when the segregations are distributed on a larger grain-boundary area. Finegrained material is therefore more ductile; the DBTT of recrystallized molybdenum is raised from - 5 0 to about +200°C when the grain grows in size from 20 to 300 jum. Purifying by means of electron-beam melting or "scavenging" with low concentrations (0.1 wt. %) of reactive metals with a high affinity towards impurities (e.g., boron, aluminium, titanium, zirconium, yttrium, and cerium) reduces or neutralizes the grain-boundary substances. At already low concentrations of oxygen in molybdenum, keeping the C/O ratio greater than two is said to completely eliminate the detrimental effects of oxygen (Kumar and Eyre, 1980), apparently by means of carbon diffusing faster to free surfaces and

preventing oxygen from segregating. Annealing impure molybdenum below 830 °C was found to favor the segregation of carbon only, and to decrease therefore the DBTT (Suzuki et al, 1981). Dispersed second phases (in alloys of type TZM, etc., or in ODS alloys) can stop cracks. Alloy softening due to substitutional elements can reduce the DBTT of refractory metals by way of decreasing the yield stress at low temperatures (Pink and Arsenault, 1979). Cold working is most efficient, since the geometry of stretched-out grains prevents grain-boundary fractures, and forces the fracture to propagate in a transcrystalline mode. However, a high degree of work hardening reduces the fracture elongation. Sometimes low deformation is sufficient, not exceeding the Liiders range, in order to reduce the DBTT below room temperature without having to cope with the detrimental effects on the fracture elongation. The DBTT of Non-Sag Structures In the heavily-worked condition, there is little difference between the ductility of undoped and doped thin tungsten wires; their DBTT is generally below room temperature (Seigle and Dickinson, 1963). In the recrystallized state, the overlapping grains prevent failure due to bending stresses. Figure 10-39 shows an example of a crack created by bending a doped molybdenum sheet with overlapping grain "flakes". It is clear that the initiation of the fracture is transcrystalline and thus more difficult. Figure 10-24 shows how far below ambient temperature the DBTT of molybdenum is situated when advantage of this effect is taken. Thin, fully recrystallized non-sag tungsten wires with high / values still have a DBTT situated well above room temperature (see Fig. 10-20).

10.7 Fundamental Mechanisms

Figure 10-39. Mainly transcrystalline fracture through elongated grains of a recrystallized non-sag molybdenum bent below its DBTT, i.e. below - 60 °C (Pink and Sedlatschek, 1969).

The Effect of Rhenium on the DBTT An addition of rhenium, iridium, or ruthenium to Via metals up to the solubility limit lowers the DBTT drastically (Wukusick, 1967). Nothing similar is observed in group Va metals; the good workability of tantalum is rather impaired by an addition of rhenium. Various explanations have been proposed (Booth et al., 1965), but a satisfactory one is still required. An initial explanation was based on the idea that rhenium scavenges the oxygen, forming complex oxides which do not wet and embrittle the boundaries. Section 10.7.2 contains the experimental evidence for an increased solubility due to rhenium alloying by way of clustering. Oxygen should be removed from dislocations and grain boundaries, but it is redistributed within the lattice, assuring a better performance (Novick and Machlin, 1968). Substitutional clustering, also said to be a consequence of alloying with sufficient amounts of rhenium (Wukusick, 1967), should affect the dislocation mobility and thus the ductility. The details of this model are as evasive as those that are based on the presence or absence of second-phase films on grain-boundary

635

surfaces. According to studies with the electron microscope, dilute rhenium additions indeed increase the mobility of screw dislocations by allowing additional slip systems to operate at low temperatures (Stephens, 1970). The rhenium ductilizing effect is not as straightforward as it usually appears; there are two minima in the curve of the DBTT versus the rhenium concentration. The second, and possibly more spectacular reduction in the DBTT at higher rhenium concentrations is the proper "rhenium effect", which has been discussed so far. The first minimum at low rhenium concentrations must be traced back to alloy softening (Pink and Arsenault, 1979; Lundberg et al., 1985). 10.7.6.2 High-Temperature Fracture of Non-Sag Structures In high-temperature creep testing of straight, non-sag tungsten wires two modes of fracture were observed (Horacsek, 1989 b): ductile, transcrystalline fracture after chisel-type necking at high stresses stemming from preceding deformation by means of dislocation creep within the grains, and brittle, intercrystalline fracture at low stresses from diffusional (grainboundary) glide due to grain-boundary cavitation. The deformation in the brittle range is small, it may be as little as 0.1 %. In incandescent lamps where non-sag wires experience very low stresses at high temperatures, the fracture is thus brittle. 10.7.6.3 Fracture of Tungsten Heavy Alloys In heavy metals, four modes of fracture are possible and occur more or less simultaneously (German et al., 1984; Edmonds, 1991): (1) ductile fracture through the matrix, (2) fracture along the tungsten-ma-

636

10 Refractory Metals and Their Alloys

trix interface, (3) fracture along the grain boundary between two adjacent tungsten grains, and (4) cleavage through the tungsten spheroids. Depending on the testing conditions and types of heavy metals, certain failures are favored. Cleavage through tungsten grains disappears when the testing temperature is raised (in accord with the general tendency of pure tungsten to deform in a ductile manner at temperatures above the DBTT), and the contribution of ductile matrix failure predominates. Grain-boundary failures at tungsten-tungsten interfaces are distributed uniformly at all temperatures because they are the weakest links. Fracture along tungsten-matrix interfaces is influenced by the rate of cooling from sinter or annealing temperatures, but the results are ambiguous; both slow and fast cooling can increase the amount of inter facial fracture. Auger spectroscopy and scanning electron microscopy revealed as possible causes for interfacial fracture, impurity segregations of sulfur, phosphorus, or silicon, as well as intermetallic segregations. The suppression of such segregations by fast cooling favors a more pronounced matrix fracture, but at the same time the matrix may be strengthened less since precipitation hardening is suppressed. 10.7.7 Structure-Related Anelastic Properties

Besides being used for solving general metal-physical problems (see Chap. 12 by De Batist in Vol. 2B of this Series), internal-friction measurements have been conducted on refractory metals in order to understand interaction processes between alloying elements as well as structural changes, and their contribution to strengthening. Only these latter effects are of interest here.

10.7.7.1 Snoek-Type Damping

The Snoek peak arises from the ordering of interstitially dissolved elements in the metal lattice when an alternating stress state is applied. Secondary peaks stem from substitutional-interstitial Snoek reactions. With the help of both peaks, scavenging or aging processes can be monitored. Precipitation starts when both peaks decrease simultaneously. Secondary peaks are easily found for group Va refractory alloys, e.g., Nb-Zr alloys (Thurber et al., 1966; Szkopiak and Ahmad, 1969). However, also for molybdenum and tungsten with their low interstitial contents (and resulting experimental difficulties), it could be proved from the observance of primary and secondary peaks that the solubility of oxygen is increased when titanium (Sung and Ma, 1966) or rhenium (Booth et al., 1965) are added. 10.7.7.2 Damping from "Particle" Formation and Dissociation

Wide peaks with damping values higher than Snoek peaks were observed for Via metals (Schnitzel, 1964). For quenched molybdenum the peak extends between 100 and 500 °C (at 1 Hz), changing its height in a very complex way upon repetition of the test runs and upon aging intervals, and also disappearing when measurements were taken at decreasing temperature (Shiwa et al., 1990). This was explained by the formation and dissolution of "particles" or interstitial agglomerations about 2 nm in diameter, which are, in contrast to normal-sized precipitates, unstable due to the contribution of the interfacial energy.

10.8 Recent Developments and Future Prospects

10.7.7.3 Grain-Boundary Damping

Above 0.4 Tm, a broad peak of high damping value occurs due to the viscous behavior of the grain boundaries. Tungsten and some of its alloys were tested in this high-temperature range (Schnitzel, 1959; Berlec, 1970; Kostkowski and Stolarz, 1976). The level of damping is proportional to the number of grain boundaries present. Measuring this damping at increasing temperature provides information on structural changes; recrystallization is recognized by a sudden drop. The differences in the recrystallization temperature of pure, doped, or thoriated tungsten and of tungsten-rhenium alloys were thus documented.

10.8 Recent Developments and Future Prospects In order to cope with the growing demands for refractory metals, the capacities and capabilities of conventional production facilities increased enormously in just a few decades. Since the contamination picked up from the processing containers increases with the temperature, containerless processing could be important. This can be done in the microgravity environment of space (sintering in space has in fact helped to avoid segregation in tungsten heavy metal). Less costly techniques (since they are earth-bound) are acoustic, aerodynamic, and electromagnetic levitation, which have been successfully tried for refractory metals (Weber et al., 1990). However, all these wondrous methods can hardly compete as processes for large-scale production. Instead, advanced P/M is today capable of achieving purities previously not considered possible. So it may not even be necessary to rely on such still costly methods as

637

multiple electron-beam melting to guarantee the low purity levels of 1 ng/g needed by the electronics industry (Anonymous, 1989; Diesburg and Castel, 1989). New production methods, which are still under development, are designed to obtain fine-grained refractory metals of high hardness and strength. In "melt spinning" the liquid metals are splash-quenched to form ribbons and filaments (Anonymous, 1988 b). Fine-grained refractory alloys could also be produced employing powder consolidation by energetic high-current pulsing (Raghunathan et al., 1991). Great progress has been achieved in finding new alloys, but the observance of metal-physical principles still promises further refinement. Dispersion hardening, which has been used for a long time, has been improved through introducing new hardening agents. It appears that optimizing the mechanical properties through age hardening is not done to the extent possible. Theoretical considerations, e.g., on how to harden by coherent precipitates, should also be applied when the alloys are produced on a commercial scale. The thermomechanical treatment of alloys should be used more frequently. The tendency to favor niobium, due to its low density and cost, over tantalum will increase. Preventing single grains from covering the entire cross section of doped tungsten filaments is still one of the unsolved problems in lamp-wire research. ODS molybdenum, ODS-ZHM, and ODS-MHC alloys, and the alloys of molybdenum and tungsten with low (3 wt. %) rhenium contents (also in ODS varieties) possessing high strength and room-temperature ductility will gain more commercial appreciation. With the introduction of oxide dispersions it should become possible to produce an elongated and interlocking, recrystallized

638

10 Refractory Metals and Their Alloys

grain structure not only in thin wires, but also in thick rods or sheets. Thoriated tungsten as a material with a low electronic work function will be replaced by tungsten containing rare-earth oxides. Tungsten heavy alloys and contact materials will become available as wrought products, formed into sheets, rods, and nearly finished shapes. Research and development of chromium-base alloys will concentrate on improving the low-temperature ductility and high-temperature strength, as well as on low- and high-temperature corrosion resistance against severely attacking gaseous and liquid media. Long-range development, financed by national or international organizations rather than by company efforts, will have to be started to utilize the technical potential of this widely unknown refractory metal, of which there is no shortage in ore resources. Composite materials, sag-resistant at high temperatures, using crystalline as well as amorphous refractory metals as the matrix or as the reinforcing component, together with sophisticated designs will greatly extend the application range of refractory metals. Attempts to find effective measures against reactions with gases will continue to be made. The still very severe problem of low oxidation resistance of all refractory metals will stimulate the development of protective coatings, including low-cost diffusional layers.

10.9 References Anonymous (1980), in: Metals Handbook, Vol. 3, 9th ed.: Metals Park, OH: ASM Int., pp. 314-329. Anonymous (1988 a), Vanadium. Albany, OR: Teledyne - Wah Chang. Anonymous (1988 b), Int. J. Refract. Hard Met. 7, 15. Anonymous (1989), Int. J. Refract. Met. Hard Mater. 8, 90-92.

Begley, R. T, Harrod, D. L., Gold, R. E. (1968), in: Refractory Metal Alloys: Metallurgy and Technology. Machlin, I., Begley, R. T., Weisert, E. D. (Eds.). New York: Plenum, pp. 41- 83. Belz, L. H. (1959), in: Reactive Metals, Metall. Soc. Conf., Vol. 2: Clough, W. R. (Ed.). New York: Interscience, pp. 525-539. Benesovsky, F. (1979), Wolfram Gmelins Handbuch der anorganischen Chemie. Erganzungsband Al, 8. Aufl. Berlin: Springer. Berlec, I. (1970), Metall. Trans. 1, 2677-2683. Bildstein, H., Eck, R. (1977), in: Preprints 9th Plansee Seminar, Vol. 1: Benesovsky, F. (Ed.). Reutte: Metallwerk Plansee, No. 10. Booth, J. G., Jaffee, R. I., Salkovitz, E. I. (1965), in: Proc. 5th Plansee Seminar: Benesovsky, F. (Ed.). Reutte: Metallwerk Plansee, pp. 547-570. Buckman, R. W., Jr. (1988), in: Alloying: Walter, J. L., Jackson, M. R., Sims, C. T. (Eds.). Metals Park, OH: ASM Int., pp. 419-445. Carlen, J. C , Bryskin, B. D. (1992), Int. J. Refract. Met. Hard Mater. 11, 343. Chandler, W. T., Walter, R. X (1968), in: Refractory Metal Alloys: Metallurgy and Technology: Machlin, I., Begley, R. X, Weisert, E. D. (Eds.). New York: Plenum, pp. 197-249. Chang, W. H. (1964), Trans. Am. Soc. Met. 57, 5237553. Christian, J. W. (1983), Metall. Trans. 14A, 12371256. Clare, J. H., Crawmer, D. E. (1982), in: Metals Handbook, Vol. 7, 9th ed., Metals Park, OH: ASM Int., pp. 361-374. Danninger, H., Lux, B. (1991), unpublished results. Diesburg, D. E., Castel, E. D. (1989), in: Proc. 12th Plansee Seminar, Vol. 3: Bildstein, H., Ortner, H. M. (Eds.). Reutte: Metallwerk Plansee, pp. 371-378. Dorn, J. E., Rajnak, S. (1964), Trans. Metall. Soc. AIME 230, 1052-1064. Eck, R. (1974), Pulvermetallurgie 22, 165-174. Eck, R. (1978), Metall 32, 891-894. Eck, R. (1979), Pulvermetallurgie 27, 53-74. Eck, R. (1980), Wire Ind. 47, 523-527. Eck, R. (1981), in: Proc. 10th Plansee Seminar, Vol. 3: Ortner, H. M. (Ed.). Reutte: Metallwerk Plansee, pp. 1-11. Eck, R. (1984), Austrian Patent 386 843, Eur. Patent 172, 852. Eck, R (1985 a), in: Proc. 11th Plansee Seminar, Vol. 2: Bildstein, H., Ortner, H. M. (Eds.). Reutte: Metallwerk Plansee, pp. 39-57. Eck, R. (1985b), Ceram. Eng. Sci. Proc. 6 (3-4), 154-166. Eck, R. (1989), in: Proc. 12th Plansee Seminar, Vol. 1: Bildstein, H., Ortner, H. M. (Eds.). Reutte: Metallwerk Plansee, pp. 1047-1063. Eck, R., Pink, E. (1992), Int. J. Refract. Met. Hard Mater. 11, 337-341.

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Kumar, A., Eyre, B. L. (1980), Proc. R. Soc. London A 370, 431-458. Lambert, B., Droegkamp, R. E. (1984), in: Metals Handbook, Vol. 7, 9th ed.: Metals Park, OH: ASM Int., pp. 765-772. Lambert, I B. (1990), in: Metals Handbook, Vol. 2, 10th ed.: Metals Park, OH: ASM Int., pp. 557585. Le Lay, L. (1990), in: Metals Handbook, Vol. 2, 10th ed.: Metals Park, OH: ASM Int., pp. 1077-1081. Leichtfried, G., Wetzel, S. (1989), in: Proc. 12th Plansee Seminar, Vol. 1: Bildstein, H., Ortner, H. M. (Eds.). Reutte: Metallwerk Plansee, pp. 10231046. Li, C. I, German, R. M. (1984), Int. J. Powder Metall. Powder Technol. 20, 149-162. Linke, X, Akiba, M., Ando, T, Coad, X P., Deschka, S., Wallura, E. (1992), Int. J. Refract. Met. Hard Mater. 11, 357-365. Lison, R., Ambroziak, A., Horn, H. (1989), in: Proc. 12th Plansee Seminar, Vol. 1: Bildstein, H., Ortner, H. M. (Eds.). Reutte: Metallwerk Plansee, pp. 7 5 105. Lundberg, L. B., Ohriner, E. K., Tuominen, S. M., Whelan, E. P. (1985), in: Physical Metallurgy and Technology of Molybdenum and Its Alloys: Miska, K. H., Semchyshen, M., Whelan, E. P. (Eds.). Greenwich, CT: AMAX, pp. 71-79. Lupton, D. (1991), Sondermetalle, Bericht Nr. 2. Ostfildern, Germany: Techn. Akademie Esslingen. Marschall, R. W, Holden, F. C. (1966), in: High Temperature Refractory Metals, Metall. Soc. Conf. Vol. 34: Fountain, R. W, Malt, X, Richardson, L. S. (Eds.). New York: Gordon & Breach, pp. 129-159. Mazey, D. X, English, C. A. (1984), /. Less-Common Met. 100, 385-427. Neges, X (1994), Diploma Thesis, Montanuniversitat Leoben. Northcott, C. (1956), Molybdenum. London: Butterworth. Novick, D. X, Machlin, E. S. (1968), Trans. Am. Soc. Met. 61, 777-783. Ogura, S., Hayashi, K. (1991), in: Tungsten 1990 (5th Int. Tungsten Symp.). Shrewsbury: MPR Publ. Serv. Ltd., pp. 140-149. Peng, K. Y, Zhao, H. (1989), in: Proc. 12th Plansee Seminar, Vol. 1: Bildstein, H., Ortner, H. M. (Eds.). Reutte: Metallwerk Plansee, pp. 547-560. Penrice, T. W. (1984), in: Metals Handbook, Vol. 7, 9th ed.: Metals Park, OH: ASM Int., pp. 688-691. Perkins, R. A. (1968), in: Refractory Metal Alloys: Metallurgy and Technology: Machlin, I., Begley, R. T., Weisert, E. D. (Eds.). New York: Plenum, pp. 85-120. Pink, E. (1986), /. Mater. Sci. 3, 450-451. Pink, E. (1986), Int. J. Refract. Hard Met. 5,192-199. Pink, E., Arsenault, R. X (1979), Prog. Mater. Sci. 24, 1-50.

Pink, E., Bartha, L., (Eds.) (1989), The Metallurgy of Doped/Non-Sag Tungsten. London: Elsevier. Pink, E., Gaal, I., (1989), in: The Metallurgy of Doped/Non-Sag Tungsten: Pink, E., Bartha, L. (Eds.). London: Elsevier, pp. 209-233. Pink, E., Sedlatschek, K. (1969), Metall 23, 12491257. Pink, E., Schrank, X, Shiwa, Y, Eck, R. (1985), in: Proc. 11th Plansee Seminar, Vol. 1: Bildstein, H., Ortner, H. M. (Eds.). Reutte: Metallwerk Plansee, pp. 211-228. Pionke, L. X, Davis, X W. (1979), in: Technical Assessment of Niobium Alloys Data Base for Fusion Reactor Application: St. Louis: McDonnell Douglas, p. 30. Predmore, R. E., Arsenault, R. X, Sparks, C. X, Jr. (1972), in: Proc. 1971 Int. Conf on the Mechanical Behavior of Materials, Vol. I. Tokyo: The Society of Materials Science, Japan, pp. 18-30. Raghunathan, S. K., Persad, C , Bourell, D. L., Marcus, H. L. (1991), Mater. Sci. Eng. A 131, 243-253. Rieck, G. D. (1967), Tungsten and Its Compounds. Oxford: Pergamon. Rostocker, W. (1958), The Metallurgy of Vanadium. New York: Wiley. Ryan, N. E., Martin, X W. (1969a), in: Proc. 6th Plansee Seminar: Benesovsky, F. (Ed.). Reutte: Metallwerk Plansee, pp. 182-207. Ryan, N. E., Martin, X W. (1969 b), J. Less-Common Met. 17, 363-376. Savitsky, Y M., Povarova, K. B., Makarov, P. V, Zavarzina, Y K. (1977), Planseeber. Pulvermetall. 25, 168-185. Savitsky, Y M., Povarova, K. B., Makarov, P. V. (1982), Z. Metallkd. 73, 92-97. Schider, S., Plenk, H., Pfluger, G., Thoma, H., Stickler, R., WeiB, B., Gartner, T., Ennemoser, K. (1985), in: Proc. 11th Plansee Seminar, Vol. 3: Bildstein, H., Ortner, H. M. (Eds.). Reutte: Metallwerk Plansee, pp. 9-36. Schnitzel, R. H. (1959), in: Reactive Metals, Metall. Soc. Conf. Vol. 2: Clough, W. R. (Ed.). New York: Interscience, pp. 245-263. Schnitzel, R. H. (1964), Trans. Metall. Soc. AIME 230, 609-611. Seigle, L. L., Dickinson, C. D. (1963), in: Refractory Metals and Alloys II, Metall. Soc. Conf. Vol. 17: Semchyshen, M., Perlmutter, I. (Eds.). New York: Gordon & Breach, pp. 65-116. Sestak, B., Seeger, A. (1978), Z. Metallkd. 69, 195202, 355-363, 425-432. Shen, Y S., Lattari, P., Gardner, X, Wiegard, H. (1990), in: Metals Handbook, Vol. 2, 10th ed.: Metals Park, OH: ASM Int., pp. 840-868. Shiwa, Y, Stiiwe, H. P., Pink, E. (1990), Ada Metall. Mater. 38, 819-824. Slade, P. G. (1986), Int. J. Refract. Hard Met. 5, 208-214. Smathers, D. B. (1990), in: Metals Handbook, Vol. 2, 10th ed.: Metals Park, OH: ASM Int., pp. 10601076.

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Smith, D. L., Loomis, B. A., Diercks, D. R. (1985), /. Nucl. Mater. 135, 125-139. Snow, D. B. (1989), in: The Metallurgy of Doped/ Non-Sag Tunsten: Pink, E., Bartha, L. (Eds.). London: Elsevier, pp. 189-202. Stephens, J. R. (1970), Metall. Trans. 1, 1293-1301. Stephens, J. R., Petrasek, D. W., Titran, R. H. (1990), Int. J. Refractory Met. Hard Mater. 9, 96-103. Stoloff, N. S. (1969), in: Fracture, Vol. VI: Liebowitz, H. (Ed.). New York: Academic, pp. 1-81. Subhash, G., Lee, Y I , Ravichandran, G. (1994), Acta Metall. Mater. 42, 319-340. Sully, A. H., Brandes, E. A. (1967), Chromium, 2nd ed.: London: Butterworths. Sung, C. Y, Ma, Y. L. (1966), Acta Metall. Sin. (JinshuXuebao) 9, 181-188. Suzuki, S., Matsui, H., Kimura, H. (1981), Mater. Sci. Eng. 47, 209-216. Szkopiak, Z. C , Ahmad, M. S. (1969), in: Proc. 6th Plansee Seminar: Benesovsky, F. (Ed.). Reutte: Metallwerk Plansee, pp. 345-356. Thurber, W. C , Distefano, X R., Kollie, T. G., Inouye, H., McCoy, H. E., Stephenson, R. L. (1966), in: High Temperature Refractory Metals, Metall. Soc. Conf. Vol. 34: Fountain, R. W., Malt, I, Richardson, L. S. (Eds.). New York: Gordon & Breach, pp. 59-77. Tietz, T. E., Wilson, J. W. (1965), Behavior and Properties of Refractory Metals. London: E. Arnold Publ. Ltd. Trefilov, V. I., Milman, Y V. (1989), in: Proc. 12th Plansee Seminar, Vol. 1: Bildstein, H., Ortner, H. M. (Eds.). Reutte: Metallwerk Plansee, pp. 107-131. Trefilov, V. I., Moiseev, V. F. (1978), Dispersnie Chastitsy b Tugoplavkikh Metallakh. Kiev: Naukova dumka. Wadsworth, J. (1983), Metall. Trans. 14 A, 285-294. Wadsworth, J., Klopp, W. D. (1985), in: Physical Metallurgy and Technology of Molybdenum and Its Alloys: Miska, K. H., Semchyshen, M., Whelan, E. P. (Eds.). Greenwich, CT: AMAX, pp. 127-133. Wang, C. T., Webster, R. T. (1982), in: Metals Handbook, Vol. 5, 9th ed.: Metals Park, OH: ASM Int., pp. 659-666. Warnes, W H. (1990), in: Metals Handbook, Vol. 2, 10th ed.: Metals Park, OH: ASM Int., pp. 10301042. Warren, R. (1989), in: The Metallurgy of Doped/NonSag Tungsten: Pink, E., Bartha, L. (Eds.). London: Elsevier, pp. 293-301. Weber, J. K. R., Krishnan, S., Nordine, P. C. (1990), in: Refractory Metals: Extraction, Processing and Applications: Liddell, K. N. C , Sadoway, D. R., Bautista, R. G. (Eds.). Warrendale, PA: TMS, pp. 169-180. Wilcox, B. A. (1968), in: Refractory Metal Alloys: Metallurgy and Technology: Machlin, I., Begley, R. T., Weisert, E. D. (Eds.). New York: Plenum, pp. 1-39.

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Wilkinson, W D. (1969), Properties of Refractory Metals. New York: Gordon & Breach. Witzke, W R., Raffo, P. L. (1971), Metall. Trans. 2, 2533-2539. Wukusick, C. S. (1967), in: Refractory Metals and Alloys IV, Vol. I; Metall. Soc. Conf. Vol. 41: Jaffee, R. I., Ault, G. M., Maltz, X, Semchyshen, M. (Eds.). New York: Gordon & Breach, pp. 231 -245. Yih, S. Y H., Wang, C. T. (1979), Tungsten: Sources, Metallurgy, Properties and Applications. New York: Plenum. Yin, X S., Lii, Z., Xing, Y. H. (1992), Int. J. Refract. Met. Hard Mater. 11, 351-356. Zitter, H., Plenk, H., Jr. (1987), J. Biomed. Mater. to. 2i, 881-896.

General Reading ASM-International Handbook Committee (Eds.) (1990, 1991, 1993), Metals Handbook, 10th ed. Metals Park, OH: ASM International. Vol. 2, pp. 565-571, 1091-1201; Vol.4, pp. 815-819; Vol. 6, 580-582,941-947. Bildstein, H., Ortner, H. M., Eck, R. (Eds.) (1985, 1989, 1993), Proc. 11th-13th Plansee Seminars. Reutte: Metallwerk Plansee. International Tungsten Industry Association (ITIA) (Eds.) (1991), Tungsten 1990 (Proc. 5th Int. Tungsten Symp.). Shrewsbury, UK: MPR Publ. Serv. Ltd. Kieffer, R., Braun, H. (1963), Vanadin, Niob und Tantal. Berlin: Springer. Kieffer, R., Jangg, G., Ettmayer, P. (1971), Sondermetalle. Vienna: Springer. Miska, K. H., Semchyshen, M., Whelan, E. P. (Eds.) (1985), Physical Metallurgy and Technology of Molybdenum and its Alloys. Greenwich, CT: AMAX. Pink, E., Bartha, L. (Eds.) (1989), The Metallurgy of Doped/Non-Sag Tungsten. London: Elsevier. Rieck, G. D. (1967), Tungsten and Its Compounds. Oxford: Pergamon Rostoker, W (1958), The Metallurgy of Vanadium. New York: Wiley. Sully, A. H., Brandes, E. A. (1967), Chromium, 2nd ed. London: Butterworth. Tietz, T. E., Wilson, X W. (1965), Behavior and Properties of Refractory Metals. London: E. Arnold. Trefilov, V. I., Mil'man, Yu. V, Firstov, S. A. (1975), Fizicheskie ocnovy prochnosti tugoplavkikh metallov. Kiev: Naukova dumka. Trefilov, V. I., Moiseev, V F. (1978), Dispersnie chastitsy b tugoplavkikh metallakh. Kiev: Naukova dumka. Yih, S. Y. H., Wang, C. T. (1979), Tungsten: Sources, Metallurgy, Properties and Applications. New York: Plenum.

11 Intermetallics Gerhard Sauthoff

Max-Planck-Institut fur Eisenforschung, Diisseldorf, Germany

List of Symbols and Abbreviations 11.1 Introduction 11.1.1 Definition of Intermetallics and Outline of This Report 11.1.2 Historical Remarks 11.2 General Considerations 11.2.1 Bonding, Crystal Structure, and Phase Stability 11.2.2 Bonding Strength and Basic Properties 11.2.3 Criteria for Phase Selection 11.3 Titanium Aluminides and Related Phases 11.3.1 Ti3Al 11.3.1.1 Basic Properties and Phase Diagram 11.3.1.2 Microstructure and Mechanical Behavior 11.3.1.3 Environmental Effects 11.3.1.4 Applications 11.3.2 TiAl 11.3.2.1 Basic Properties and Phase Diagram 11.3.2.2 Microstructure and Mechanical Behavior 11.3.2.3 Environmental Effects 11.3.2.4 Applications 11.3.3 Al3Ti and Other D0 2 2 Phases 11.3.3.1 Basic Properties and Phase Diagram 11.3.3.2 Microstructure and Mechanical Behavior 11.3.4 Trialuminides with the L l 2 Structure •11.3.4.1 Basic Properties and Phase Diagrams 11.3.4.2 Microstructure and Mechanical Behavior 11.4 Nickel Aluminides and Related Phases 11.4.1 Ni3Al 11.4.1.1 Basic Properties and Phase Diagram 11.4.1.2 Microstructure and Mechanical Behavior 11.4.1.3 Environmental Effects 11.4.1.4 Applications 11.4.2 Other Ll 2 Phases 11.4.2.1 General Remarks 11.4.2.2 LI 2 Phases of Particular Interest 11.4.3 NiAl Materials Science and Technology Copyright © WILEY-VCH Verlag GmbH & Co KGaA. Allrightsreserved.

646 647 647 648 651 651 654 657 660 660 660 662 665 667 668 668 669 673 675 676 676 677 682 682 683 684 684 684 686 690 691 693 693 694 697

644

11.4.3.1 11.4.3.2 11.4.3.3 11.4.3.4 11.4.3.5 11.4.3.6 11.4.4 11.4.4.1 11.4.4.2 11.4.4.3 11.4.4.4 11.4.5 11.4.6 11.5 11.5.1 11.5.2 11.5.3 11.6 11.6.1 11.6.2 11.6.3 11.6.4 11.6.5 11.7 11.7.1 11.7.2 11.7.3 11.7.4 11.7.5 11.7.6 11.7.7 11.8 11.8.1 11.8.2 11.8.2.1 11.8.2.2 11.8.2.3 11.8.2.4 11.9 11.10 11.10.1 11.10.2 11.11 11.11.1 11.11.2

11 Intermetallics

Basic Properties Phase Diagram and Martensitic Transformation Microstructure and Mechanical Behavior Creep Environmental Effects Alloy Developments and Applications Other B2 Phases CoAl NiTi FeTi, CoTi, CoZr, and CoHf FeCo Heusler-Type Phases Nickel-Molybdenum Phases Iron Aluminides and Related Phases Fe3Al Fe 3 AlC x and Related Phases FeAl Cu-Base Phases CuZn Cu-Zn-Al Shape Memory Alloys Cu-Al-Ni Shape Memory Alloys Cu-Au Phases Cu Amalgams A15 Phases Basic Properties V3Si V 3 Ga Nb 3 Sn Nb3Al Nb3Si Cr3Si Laves Phases Basic Properties Applications Superconducting Materials Magnetic Materials Hydrogen Storage Materials Structural Alloys Beryllides Rare-Earth Compounds Magnet Materials Hydrogen Storage Materials Silicides M3Si Phases M2Si Phases

697 699 701 704 712 713 723 724 725 726 727 729 729 730 730 732 733 736 736 737 738 739 740 740 740 741 743 743 744 746 746 746 746 748 748 749 749 750 752 753 754 754 755 756 757

Contents

11.11.3 11.11.4 11.11.5 11.12 11.13 11.14

M 5 Si 3 Phases MSi Phases Disiliddes Prospects Acknowledgements References

645

758 759 760 763 765 766

646

11 Intermetallics

List of Symbols and Abbreviations A ^4diff a, b, c B b D d G H H Hform Hfus ^totai ifvap k Xlc Ms S T Tm

dimensionless factor dimensionless factor lattice parameters magnetic induction length of Burgers vector effective diffusion coefficient effective diffusion length Gibbs energy; shear modulus enthalpy magnetic field strength formation enthalpy enthalpy of fusion total phase enthalpy enthalpy of vaporization Boltzmann's constant fracture toughness martensite formation temperature entropy absolute temperature melting temperature

s ediff a a th Q

secondary strain rate diffusion creep rate applied stress threshold stress atomic volume

APB BDTT EL HPSN LPS NASP ODS PST REPM tcp TMP UTS VEC VLSI YS

antiphase boundary brittle-to-ductile transition temperature tensile elongation hot-pressed silicon nitride long-period superlattice National Aero-Space Plane oxide-dispersion strengthened polysynthetically twinned rare earth permanent magnet topologically close-packed thermal mechanical processing ultimate tensile strength valence electron concentration very-large-scale-integrated circuit yield strength

11.1 Introduction

,,Die Sprodigkeit der metallischen Kristallarten singularer Zusammensetzung, der Metallverbindungen, kann ihren Grund darin haben, daB ihnen die Fahigkeit, Gleitebenen zu bilden, abgeht, oder daB neben dieser Fahigkeit noch die Eigentiimlichkeit der RiBbildung auftritt." ('The reason for the brittleness of metallic crystals with unique compositions, i.e. the metal compounds, may be that they lack the ability to form glide planes, or that in addition to this ability, there exists the characteristic of crack formation.") Tammann and Dahl (1923)

11.1 Introduction 11.1.1 Definition of Intermetallics and Outline of This Report

Intermetallics is the short and summarizing designation for the intermetallic phases and compounds which result from the combination of various metals, and which form a tremendously numerous and manifold class of materials, as will become clear in the following sections. According to a simple definition (Schulze, 1967; Girgis, 1983), intermetallics are compounds of metals whose crystal structures are different from those of the constituent metals, and thus intermetallic phases and ordered alloys are included. During the last ten years intermetallics have been of enormous, and still increasing, interest in materials science and technology with respect to applications at high temperatures, and a new class of structural materials is expected to be developed on the basis of intermetallics. Various materials developments are under way in various parts of the world, in particular in the U.S.A., in Japan, and in Germany. These activities are

647

the main subject of the present report. However, it has to be noted that besides intermetallics with oustanding high-temperature properties there are other intermetallics with outstanding physical properties which have given rise to developments of new functional materials much earlier, and which will also be addressed. The numerous and manifold activities on intermetallics, as well as the enormous interest, have given rise to many conferences, symposia and workshops on materials research and development of intermetallics during previous years. The results of these activities have been published in the respective proceedings volumes, as well as the publications in scientific journals and research reports, and there are an increasing number of review papers on various aspects of intermetallics in such proceedings and in the scientific journals. Furthermore, particular aspects of intermetallics are treated in various volumes of the Series Materials Science and Technology (MST): chapters by Pettifor, and Ferro and Saccone in Volume 1, chapters by Delaey, and Inden and Pitsch in Volume 5, and chapters by Umakoshi and Mukherjee in Volume 6. It is the aim of the present report to give a summarizing description of the various intermetallics which have already been selected for materials developments, or which have been, and still are, regarded as promising for materials developments. The characteristic properties are shown, the physical mechanisms which control the observed behavior and the problems involved are discussed, and the present state and prospects of the respective materials developments are assessed. This report, however, cannot be exhaustive because of the large and still increasing number of activities, and because many developments have not yet been described in the open literature.

648

11 Intermetallics

After an introductory section with a brief account of the history of intermetallics, the fundamentals of intermetallics are discussed by making reference to the respective chapters in various volumes of MST. In particular, the stability of phases, as well as the relation between atomic bonding and basic properties are addressed, and the criteria for grouping the various phases are discussed. Then the various intermetallics and the respective materials developments are described, and finally the prospects, as well as the needs for research and development, are assessed. 11.1.2 Historical Remarks The history of intermetallics has been outlined repeatedly in some detail by Westbrook, one of the central people in the research and development of intermetallics in the second half of this century (Westbrook, 1967, 1970, 1977, 1993). Therefore only some important points are noted here. Intermetallics have been made use of since the beginning of metallurgy, as is exemplified in Table 11-1 (for nomenclature see the chapter by Ferro and Saccone in Volume 1 of MST). Such intermetallics resulted from the used alloy systems with low melting temperatures, and the applications relied on the outstanding hardness and wear resistance of the intermetallics together with their metallic properties (Westbrook, 1977). Because of the metallic behavior of intermetallics they could be polished, and thus the decorative aspect was also important for many applications. Examples are the bronze coatings in ancient Egypt and the mirrors which were used, not only by the ancient Chinese, but also by Etruscans and Romans (Westbrook, 1977). In this context it has to be

mentioned that some intermetallics show beautiful colors, e.g. the violet Cu2Sb ("Regulus of Venus"), which was also discovered early on (Westbrook, 1977). Intermetallics became a subject of scientific research during the last century with the development of physical metallurgy, and the first validated instance of intermetallic compound formation was reported by Karsten (1839) in Germany (Westbrook, 1967, 1970). Intensive and broadscale work on intermetallics was initiated by Tammann in Gottingen, Germany and Kurnakov in St. Petersburg, Russia at the turn of the century, from which a large number of directive papers resulted (e.g. Kurnakov, 1900, 1914; Tammann, 1903, 1906; Kurnakov and Zhemchuzhny, 1908; Tammann and Dahl, 1923). The early work on intermetallics during the first decades of this century included studies of phase stabilities, phase equilibria and phase reactions in order to establish phase diagrams, as well as studies of the various properties, i.e. chemical and electrochemical properties, physical properties including magnetism and superconductivity, and mechanical properties, as are summarized in Tammann (1932). With respect to mechanical behavior it was realized that the outstanding hardness of intermetallics is accompanied by an unusual brittleness, and the reasons for this were also investigated (Tammann and Dahl, 1923). Intermetallics were used in this century first and primarily for applications as functional materials, as is exemplified in Table 11-1. Indeed the first industrial applications relied on the special magnetic behavior of certain phases, and respective materials developments led e.g. to Sendust, which shows outstanding magnetic properties and wear resistance and is widely used for magnetic heads in tape recorders (Yamamoto, 1980; Brock, 1986). In the

11.1 Introduction

649

Table 11-1. Some past and present applications of intermetallics (Sauthoff, 1989). Since approx. 2500 B.C. 100 B.C.

Material or process

Phase

Application

Reference

cementation

Cu3As

coating of bronze tools, etc. (Egypt, Anatolia, Britain)

Westbrook (1977) Gmelin-Institut (1955) Westbrook (1977)

yellow brass

CuZn

coins, ornamental parts (Rome)

high tin bronze

Cu 31 Sn 8

mirror (China)

600

amalgam

Ag 2 Hg 3 + Sn6Hg

dental restorative (China)

Westbrook (1977) Waterstrat (1990)

1500

amalgam

Cu 4 Hg 3

dental restorative (Germany)

Paufler (1976), Westbrook (1977) Waterstrat (1990)

1505

amalgam

Sn8Hg

mirror surface (Venice)

Westbrook (1977)

1540

type metal

SbSn

printing

Westbrook (1977)

1910

Acutal

(Cu,Mn)3Al

fruit knife (Germany)

Heusler (1989)

0

1921

Permalloy

Ni 3 Fe

high permeability magnetic alloy

Bozorth (1951)

1926

Permendur

FeCo(-2V)

soft magnetic alloy

Bozorth (1951), Chen (1961)

1931

Alnico

NiAl-Fe-Co

permanent magnet material

De Vos (1969)

1935

Sendust

Fe3(Si,Al)

magnetic head material

Yamamoto (1980)

1938

Cu-Zn-Al Cu-Al-Ni

CuZn-Al (Cu,Ni)3Al

shape memory alloys

Hodgson (1990)

1950

pack aluminide coating

NiAl, CoAl

surface coating for protection from environment

Nicholls and Stephenson (1991)

1956

Kanthal Super, Mosilit

MoSi2

electric heating elements

Fitzer and Rubisch (1958)

1961

A15 compound

Nb 3 Sn

superconductors

Westbrook (1977), Geballe and Hulm (1986)

1962

Nitinol

NiTi

shape memory alloy

Delaey etal. (1974) Hodgson (1990)

1967

Co-Sm magnets

Co5Sm

permanent magnets

Stadelmaier et al. (1991), Westbrook (1977), Buschow (1986)

second half of this century another important application resulted from the development of new superconducting materials based on the A15 compounds which are used for superconducting magnets (B.W.Roberts, 1967; Westbrook, 1977; Dew-Hughes, 1986). A third group of functional materials, the shape memory alloys, makes use of a martensitic phase

transformation and has again found manifold applications during the last three decades, e.g. Otsuka and Shimizu (1986). An important group of functional materials is formed by the III-V compounds, e.g. InSb, InAs, GaAs, which have found applications in electronics and in thermoelectric power generation (C. S. Roberts, 1967; Cadoff, 1967). The constitutive ele-

650

11 Intermetallics

ments of these phases represent the transition from metals to semi-metals and nonmetals, and the resulting compounds are semiconductors. Thus these intermetallic semiconductor compounds are outside the field of intermetallics proper, which exhibit metallic behavior. Intermetallics did not find applications as structural materials in the past because of their brittleness. The only noteworthy exception is the continuing use of amalgams as dental restoratives (Westbrook, 1974, 1977). On the other hand, various intermetallics were successfully used as strengthening second phases in conventional alloys for structural applications (Westbrook, 1970). Thus it was clear that intermetallics are promising for applications as structural materials at high temperatures because of their high hardness and stability, and indeed manifold activities were started at the beginning of the fifties for clarifying the potential of intermetallics for structural applications (Westbrook, 1960 a, b). Various promising candidate phases were identified, but the respective materials developments were hampered by the unsolved brittleness problems, and thus the various activities slowly died away during the 1960s (Westbrook, 1965, 1970, 1977; Ryba, 1967; Liu and Stiegler, 1984; Cahn, 1989; Liu et al, 1990). One offspring of these activities was the development of the electric heating elements based on MoSi2 (see Table 11-1), which rely on the advantageous chemical behavior of this phase, i.e. its high oxidation resistance at very high temperatures (Fitzer and Rubisch, 1958; Schrewelius and Magnusson, 1966). In this context the surface coatings should be mentioned which have been developed in the course of superalloy developments to protect the coated materials against high temperature

corrosion, and which rely primarily on such intermetallics as NiAl and CoAl (Nicholls and Stephenson, 1991). It should be noted that the mentioned molybdenum disilicide is a borderline case of the intermetallics since silicon is not a metal, but a semiconductor. It is known that the combination of silicon with metals gives rise to compounds with metallic properties, e.g. MoSi2, as well as compounds with semiconductor properties (Nowotny, 1963), i.e., silicides mark the transition from intermetallics to compounds of metals and nonmetals. Nevertheless, silicides are traditionally included in the field of intermetallics because of their many similarities with metals (Wehrmann, 1967). The first intensive and successful structural materials developments were based on the titanium aluminides Ti3Al and TiAl, and were started at the beginning of the 1970s with fundamental deformation studies (Shechtman et al., 1974; Fleischer etal., 1989a; Schneibel etal., 1986), whereas the potential of these phases for high-temperature applications had already been recognized during the 1950s (McAndrew and Kessler, 1956). These developments are still continuing with a multitude of activities. Another similar intensive and successful structural materials development was based on the nickel aluminide Ni 3 Al, which is the strengthening second phase in the superalloys, and which was also a candidate phase during the 1950s (Liu and Stiegler, 1984). This development had an enormous impact due to the detection of the "ductilization" effect of small additions of boron (Aoki and Izumi, 1979; Liu and Koch, 1983), since this effect gave rise to large research programs on intermetallics and excited the general and still increasing interest in intermetallics. More

11.2 General Considerations

and more research groups started work on intermetallics, and in most cases during the last decade Ni3Al was selected as the subject of research. A further related development, which was already started at the beginning of the 1970s, was based on the quaternary phase (Fe,Co,Ni) 3 V with the crystal structure of Ni 3 Al (Liu and Inouye, 1979). These successful developments have initiated work on other less-common phases recently which aims at low densities and/or high application temperatures (e.g. Sauthoff, 1989). From these various developments over the years the manifold spectrum of present activities has evolved which is the subject of the subsequent sections.

11.2 General Considerations 11.2.1 Bonding, Crystal Structure, and Phase Stability

Intermetallics form because the strength of bonding between the respective unlike atoms is larger than that between like atoms. Accordingly, intermetallics form particular crystal structures with ordered

651

atom distributions where atoms are preferentially surrounded by unlike atoms. Simple examples of such crystal structures are shown in Fig. 11-1. The crystal structures of intermetallics have been discussed in detail by Ferro and Saccone in Volume 1 of MST. The crystal structure of an intermetallic is determined by the strength and character of bonding in the crystal, which depends on the particular electronic configuration. The relation between structure type and atomic properties of the constituent atoms, however, is not a simple one and thus various criteria have been used for correlating structure type and phase type, i.e., for predicting the crystal structure for a given phase or phase group. Likewise, all the intermetallics cannot be expected to show metallic bonding similar to the constituent metals. This has been discussed in detail repeatedly, and the intermetallics have been grouped traditionally according to various criteria (e.g. Nevitt, 1963; Schulze, 1967; Schubert, 1967; Laves, 1967; Girgis, 1983; Hafner, 1987; Pettifor, 1988; and the chapters by Pettifor and Ferro and Saccone in Volume 1 of MST).

TiAl

OA • B Fe3AIC

Figure 11-1. Some simple intermetallic crystal structures which are derived from the b.c.c. and f.c.c. structures, respectively, with typical examples (Sauthoff, 1989).

652

11 Intermetallics

It should be noted that these groupings are not unambiguous since different authors stress different aspects of the particular intermetallics. In the following, a brief survey on three important and rather different groups of intermetallics is given to illustrate this. The Zintl phases are formed by metals on the left-hand side and on the right-hand side of the periodic table of elements, and are characterized by completely filled electronic orbitals, normally by a full octet shell (Laves, 1967; Schafer e t a l , 1987; Schmidt, 1987). Thus they may be regarded as valence compounds which satisfy the familiar chemical valency rules (Girgis, 1983; Hafner, 1987). The Zintl phases have crystal structures which are characteristic for typical salts, e.g. NaTl with a cubic B32 structure (Strukturbericht designations are used here and in the following for the crystal structures) or Mg2Si with a cubic Cl structure (Villars and Calvert, 1991), and therefore ionic bonding is expected. However, all types of bonding (ionic, metallic and covalent) and mixtures thereof have been found depending on the particular electronic distribution, i.e. the Zintl phases may also be regarded as electronic compounds for which the electron band-structure energy is a large fraction of the total energy, and crystal structure types are related to particular valence electron concentrations (average number of valence electrons per atom) (Hafner, 1987). The best-known electron compounds are the Hume-Rothery phases which have the cubic B2 structure, e.g. p-brass at high temperatures and the transition-metal aluminides FeAl, NiAl and CoAl, or the complex cubic A13 structure ((3-manganesetype), e.g. Zn 3 Co, Cu5Si, or the closepacked hexagonal A3 structure, e.g. Cu 3 Ga, Ag3Al, for the valence electron concentration VEC = 3/2, the complex

cubic D8 2 structure (y-brass type) for VEC = 21/13, e.g. Cu 5 Zn 8 , Fe 5 Zn 2 , and again the A3 structure for VEC = 7/4, e.g. CuZn 3 , Ag 5 Al 3 (Girgis, 1983). The bonding in such intermetallics is not purely metallic in spite of the metal-like band structure. For example, the bonding in the above-mentioned NiAl has been found to be essentially covalent with some metallic character and no ionic component (Fox and Tabbernor, 1991), which can be understood in view of the band structures of Ni, Al and NiAl (Engell etal., 1991). It has indeed been shown on the basis of recent ab initio calculations that - strictly speaking - the B2 aluminides are not HumeRothery electron compounds (Schultz and Davenport, 1993). This means that the Hume-Rothery rules, which relate crystal structure to valence electron concentration, represent a very simplified picture of the bonding situation. Nevertheless, such rules are useful since they describe reality surprisingly well. Other intermetallics are best characterized by the size ratios and packing schemes of the constituent atoms, and are therefore called size-factor compounds or topologically close-packed intermetallics or Frank-Kaspar phases (Wernick, 1967; Girgis, 1983; Watson and Bennett, 1985; Hafner, 1987). One important group of these intermetallics are the A15 phases, which were mentioned in the preceding section. The best-known size-factor compounds are the Laves phases, which form the most numerous group of intermetallics and which crystallize in the closely related hexagonal C14, cubic C15 or hexagonal C36 structures with MgZn 2 , MgCu 2 , MgNi 2 , respectively, as typical examples (Wernick, 1967). The valence electron concentration, however, is also an important factor here since it decides between the three crystal structures. These crystal

11.2 General Considerations

structures are characterized by high symmetry, high atom coordination and high density in close analogy to metals, and indeed the Laves phases primarily show metallic bonding according to the limited information available (Schulze, 1967; Hafner, 1987). There are phases which form by a phase transition with a lowering of the crystal symmetry from the solid solution of the constitutent elements at a transition temperature below the melting temperature, and which are known as Kurnakov phases (Kornilov, 1967). In the simplest case the phase transition results from an ordering reaction in the solid solution lattice from which a superlattice structure is established. Examples of phase formation by atomic ordering are the phases Fe3Al with a DO3 structure, which forms by two-step ordering from the b.c.c. solid solution (see Sec. 11.5.1), as well as Ni 3 Fe with an L l 2 structure (see Sec. 11.4.2.2), CuAu with an L l 0 structure and Cu 3 Au with an L l 2 structure (see Sec. 11.6.4), the latter three of which form from f.c.c. solid solutions. This atomic ordering with superlattice formation obviously results from a stronger bonding between unlike atoms compared with bonding to like atoms. The difference in bonding energy between unlike and like atoms is usually known as the interaction energy and corresponds to an ordering energy. Thus a higher crystal stability with a higher ordering temperature is expected for a higher interaction energy. The degree of order, which depends on temperature and composition, as well as the order/disorder temperature can be modeled as a function of the interaction energy for a given phase to various degrees of approximation, e.g. to the first approximation by the simple Bragg-Williams model or to higher approximations by the cluster variation method or the Monte Carlo method

653

(De Fontaine, 1979; see also the chapter by Binder, and Inden and Pitsch in Volume 5 of MST). Apart from the degree of approximation, the predictive power of such models with respect to phase stability and phase diagram calculations relies on the knowledge of the interaction energy, and on whether such interaction energies are sufficient for the characterization of the atomic bonding state. Analyses of various transition metal systems have shown that modeling with the cluster variation method or Monte Carlo method gives good agreement with the experimental observations in most cases, whereas in some cases with b.c.c. structures the simple Bragg-Williams model gave better results than the more sophisticated models (chapter by Inden and Pitsch in Volume 5 of MST). It follows that single simple atomic parameters like the interaction energy, atomic size or valence electron concentration have only a limited validity and are generally not sufficient for predicting the phase stability as a function of the constituent elements. A better and more general correlation between the atomic properties and the compound crystal structure should be obtained if the particularities of the electronic configurations of the constituent atoms are taken into account, which are reflected by the positions in the periodic table of elements. A better correlation has indeed been found, but the resulting structure maps still have their limitations because the changes in electronic configuration due to compound formation are not considered (Pettifor, 1988; Villars et al., 1989). It has to be concluded that the bonding character and crystal structure of an intermetallic can be predicted in a significant way only on the basis of a quantum-mechanical, ab initio calculation for the re-

654

11 Intermetallics

spective phase (Hafner, 1987, 1989; Majewski and Vogl, 1989). Much progress has been made in this respect (Gyorffy et al., 1991), and for various important phases with not too complicated structures the crystal energy has been calculated as a function of the lattice structure in order to determine the phase stability (Freeman etal., 1991; Bose etal., 1991; Lin etal., 1991; Pettifor and Aoki, 1991; Gonis et al., 1991; Becker etal., 1991; Vignoul etal., 1991; Sluiter etal., 1990; Huang etal., 1991; De Fontaine etal., 1991; Turchi et al., 1991); see also the chapter by Pettifor in Volume 1 of MST. However, these calculations are very time consuming, even in simple cases, and more progress is necessary in the future in order to give guidance to practical materials developments on the basis of multinary (i.e. ternary, quaternary, etc.) phases with less simple crystal structures. From the discussion in this section it follows that the intermetallics do not form a homogeneous group of materials at all. Instead the term intermetallics comprises a huge variety of phases which differ drastically with respect to bonding, crystal structure and properties. Thus the properties of intermetallics cannot be discussed generally with respect to all intermetallics; they can only be discussed by referring to specific groups of intermetallics. It may be supposed that the particular crystal structure of an intermetallic reflects the character and strength of bonding in a sensible way, and thus the crystal structure would be a good criterion for phase classification. However, this does not mean that intermetallics with the same crystal structure are similar with respect to bonding and properties since, e.g. the B2 structure is common to the intermetallic NiAl and to the purely ionic salt CsCl, and in the case of the transition-metal disilicides with a

hexagonal C40 structure, NbSi 2 is metallic whereas CrSi2 is a semiconductor (Nowotny, 1963). It has indeed been shown recently on the basis of ab initio calculations that atomic coordination and connectivity are more relevant to crystal energy and stability than symmetry (Shah and Pettifor, 1993). In view of the complexity of the classification of intermetallics, intermetallics are often grouped according to more practical criteria which refer to similarities in behavior. In the subsequent sections the intermetallics to be discussed are grouped according to constitution and crystal structure in a somewhat subjective manner. 11.2.2 Bonding Strength and Basic Properties A basic property is the melting temperature since it is known that materials' parameters which characterize the deformation behavior are well correlated with the melting temperature (Frost and Ashby, 1982). Examples are the elastic moduli which not only control the elastic deformation, but are also important parameters for describing the plastic deformation, and the diffusion coefficients which control not only the kinetics of phase reactions, but also the kinetics of high-temperature deformation, i.e. creep. Furthermore, the melting temperature is intuitively regarded as a measure of the phase stability since it limits the application temperature range. However, a phase melts when the Gibbs energy of the solid phase, which is a thermodynamic state function and which controls the phase stability, is higher than that of the liquid phase [see e.g. Denbigh (1971) and chapter by Pelton in Volume 5 of MST]. The Gibbs energy G of a phase is given by G = H-TS

(11-1)

11.2 General Considerations

i.e., it is given by the enthalpy H at a temperature T = 0 K, and its temperature dependence is determined by the entropy S. Thus the melting temperature depends on the differences in enthalpy and entropy between the solid and liquid states and is a rather complex function of the bonding strength. Only the enthalpy is directly related to the internal crystal energy which is determined by the bonding strength. Nevertheless, the melting temperature is correlated surprisingly well with the phase formation enthalpy for sufficiently similar phases, as is shown in Fig. 11-2. For the small number of phases concerned, there is a linear relationship for the phases with a common B2 structure, whereas the phase TiAl with a different crystal structure follows another relationship. It may be concluded that the phase formation enthalpy may be a better parameter for characterizing bonding strength and phase stability, and for correlating this with the basic properties, e.g. elastic moduli. Formation enthalpies have been determined experimentally (Hultgren, 1963),

655

and they can be calculated by quantummechanical calculations (e.g., Hackenbracht and Kiibler, 1980), or estimated by more or less empirical approaches (de Boer et al., 1988). Figure 11-3 shows some data for the Ni-Al system which illustrate the good agreement between theory, estimate and experiment. However, the enthalpy values in Figs. 11-2 and 11-3 refer to compound formation from the constituent elements at room temperature, i.e. from the solid element crystals, and thus these formation enthalpies only characterize the bonding strength changes during compound formation. For complete characterization of the bonding strength the total phase enthalpy is required, which refers to the compound formation from the elements in the gaseous state. The total phase enthalpy Hioial of a phase C may be estimated by adding up the formation enthalpy Hform and the enthalpies of fusion and vaporization // fus and // vap of the constituent elements A and B, according to the reaction cycle given in Fig. 11-4. As an example, Fig. 11-5 shows

2000 -1500 ' 1000

2 Q.

10

20

30

40

50

Figure 11-2. Melting temperature as a function of formation enthalpy (Hultgren, 1963) for the B2 phases CuZn ( + ), FeAl (#), CoAl (x), NiAl (*) and for the Ll 0 phase TiAl (o) (Engell et al., 1991).

60

formation enthalpy in kj/mol of atoms

Figure 11-3. Formation enthalpies from experiments ( x ) (Hultgren, 1963), quantum-mechanical calculations ( + ) (Hackenbracht and Kiibler, 1980) and estimates (o) according to de Boer et al. (1988), as a function of composition for the Ni-Al system (Engell et al., 1991).

656

11 Intermetallics

Figure 11-4. Reaction cycle for the formation of a phase C from the gaseous elements A and B (see text).

data for the Ni-Al system, and again the agreement between estimates and experimental data is visible. It is noted that the formation enthalpies are small compared with the total enthalpies, i.e. the enthalpies of the constituent elements are the major contributions to the total enthalpies. The correlation between the total phase enthalpy and the Young's modulus, which is one of the three elastic moduli, is shown in Fig. 11-6 for various cubic phases, i.e. for the f.c.c. elements Al and Ni (Al), the f.c.c.-ordered Ni 3 Al (Ll 2 ), the b.c.c. ordered FeAl, NiAl and CoAl (B2), and the cubic Laves phases CaAl 2 , YA12, LaAl 2 , NbCr 2 , ZrCo 2 and HfCo2 (C15). The correlation is quite good for the Laves phases, i.e. there seems to be a common scatter band and the data of some other phases are near this scatter band. It has to be

noted, however, that the data points may be subject to appreciable experimental uncertainties, as is indicated by the two data points for Ni 3 Al. Such uncertainties for the Young's modulus may be due to differences in the microstructure, i.e. texture and artefacts. Apart from this, the data for the B2 phases show that the Young's modulus is not a simple function of the phase enthalpy. This has to be expected because the Young's modulus is a complex function of the potential wells of the atoms, which cannot be characterized by a single parameter. A similar correlation is found for the activation energy of diffusion, as is illustrated in Fig. 11-7. It has to be concluded that there are physically justified correlations between parameters which characterize the bonding strength and the phase stability on the one hand, and the macroscopic phase behavior on the other. However, such correlations represent complex functional relationships, and thus they are useful only for order-of-magnitude predictions. More quantitative predictions have to consider the character and strength of bonding in a more detailed way, i.e. they have to rely on quantum-mechanical calculations which are cumbersome and time consuming. Recently basic properties - in particular the elastic moduli - of various simple phases have been studied by ab initio cal-

-300

"3

Figure 11-5. Total phase enthalpy as a function of composition for the Ni-Al system: experimental data ( + ) (Hultgren, 1963) and estimates (x) according to de Boer et al. (1988) compared with the ideal solid solution (*) with vanishing formation enthalpy (Von Keitz and Sauthoff, 1991).

E

-350-

: -400-

-450 0.4 0.6 Al content

11.2 General Considerations a 400 x -

+

Ni 3 Al

ACoAl

300

FeAl

I/) ZD

A

NbCr2

ANiAl

Ni

^

^

ZrCo 2 LoAl2

CQA12

100

^

YAl2 -""Al

0 -300

-400

-500

-600

657

cleavage strength (Fu and Yoo, 1992 a). However, it has also been shown for B2 transition-metal aluminides that the ductility or brittleness is too complex a function of the strength and character of bonding to establish simple relations between the electronic distribution parameters and the mechanical properties (Schultz and Davenport, 1992, 1993).

total phase enthalpy in kJ/mol of atoms

11.2.3 Criteria for Phase Selection

(Hfotal' T=298K

Figure 11-6. Young's modulus as a function of total phase enthalpy (Hultgren, 1963; Engell et al., 1991; Von Keitz and Sauthoff, 1991) at room temperature for the f.c.c. elements Al and Ni (Al) (A) (Lide, 1992), the f.c.c.-ordered Ni3Al with an Ll 2 structure ( + ) (Davies and Stoloff, 1965; Munroe and Baker, 1988), the b.c.c. ordered FeAl, NiAl and CoAl with a B2 structure (A) (Harmouche and Wolfenden, 1985), and the cubic Laves phases CaAl2, YA12, LaAl2, NbCr 2 , ZrCo 2 and HfCo2 with a C15 structure (x) (Shannette and Smith, 1969; Schiltz and Smith, 1974; Fleischer et. al., 1988).

culations, and a lot of progress has been made in this field (e.g. Fu and Yoo, 1990, 1991; Lee and Yoo, 1990; Guo et al., 1991; Yoo and Fu, 1991, 1993; Yoo, 1991; Fu, 1991). Thus a basic understanding of the bonding character of, e.g. transition-metal aluminides has been reached and it has been shown that the degree of directional d-electron bonding determines the ideal

Ni 3 Al

A

300-

A A

c

aj

200-

>

100-

A

Candidate phases for materials developments are selected in view of specific applications, i.e. they must show certain properties which make them promising for a particular application. For permanent magnets the "energy product" \B H\ is a figure of merit and should reach high values (e.g. Stadelmaier et al., 1991), whereas for other functional materials other physical parameters are decisive. Selection criteria with respect to structural applications at high temperatures have already been discussed by Sauthoff (1989). First, such phases must have sufficient strength at the service temperature which also means sufficient creep resistance. The creep resistance scales with the diffusion coefficient and with the shear modulus (e.g. Jung et al., 1987), and both parameters scale with the melting temperature (Frost and Ashby, 1982). Thus the

NiAl

Ni

Figure 11-7. Activation energy of diffusion as a function of total phase enthalpy for Ni, Al, NiAl and Ni3Al (Frost and Ashby, 1982; Wever et al., 1989; Engell et al., 1991; Von Keitz and Sauthoff, 1991).

Al

a 1 0-250 -300

1

-350

-400 -450 -500

-550

-600

total phase enthalpy in kJ/mol of atoms

658

11 Intermetallics

candidate phase should have a sufficiently high melting temperature. The limiting temperature for structural applications of conventional metallic alloys is of the order of 75 % of the melting temperature in many cases. The most advanced, high-temperature alloys are the Ni-base superalloys, and they are used in gas turbines with service temperatures of up to about 1100°C (Petrasek etal., 1986). Thus phases with melting temperatures above 1600°C have to be considered if application temperatures above those of superalloys are aimed at. High-melting phases may form low-melting eutectics which may affect the phase stability during processing or service, and therefore such phases with low-melting eutectics should be avoided. For many functional and structural applications density is a very important and sometimes decisive parameter. Structural intermetallics for moving components must have sufficiently high specific strength, which is the ratio of strength and weight density, and which has the dimension of length. The specific fracture strength, i.e. the rupture length, of advanced superalloys, e.g. the mechanically alloyed, oxide-dispersion strengthened (ODS) superalloy MA 6000, is of the order of 15 km (Inco, 1982), which has to be surpassed by new structural materials. Phases which contain light elements, e.g. Ti, Al, Si, or Mg, may have such a low density that they compare favorably with conventional alloys in spite of lower strength or a restricted temperature range. The main problem with strong intermetallics is their brittleness, which makes processing and application difficult or impossible. However, the brittleness of intermetallics should be less severe than that of ceramics because the atomic bonding of intermetallics is, at least partially, still metallic, whereas it is primarily covalent or

H00

0

0.1 0.2 0.3 0.4 0.5 0.6 0.7

0.8 0.9

1.0

T/T m

Figure 11-8. Yield stress and deformability as a function of homologous temperature T/Tm (with T^ = melting temperature) for various materials with different types of atomic bonding (Westbrook, 1965).

ionic in ceramics. This is illustrated in Fig. 11-8 which compares the temperature dependence of the yield stress and deformability of the intermetallic NiAl with that of nickel, and of silicon and alumina. NiAl, which is a candidate phase for hightemperature applications, softens with rising temperature at about half the melting temperature Tm, which is the range for Si with covalent bonding and which is higher than that of the ceramic A12O3. However, the brittle-to-ductile transition temperature is about 0.4 Tm for NiAl, whereas it is above 0.8 Tm for Si and A12O3, i.e. the deformation behavior of such an intermetallic may be regarded as intermediate between metals and ceramics. Plastic deformation is more difficult in intermetallics than in metals and conventional metals alloys because of the stronger atomic bonding and the resulting ordered atomic distribution, which gives rise to more complex crystal structures (Paufler, 1985). Experience indeed shows that the brittleness of intermetallics increases with decreasing lattice symmetry and increasing

11.2 General Considerations

unit cell size (Paufler, 1976). Therefore intermetallics with high crystal symmetry possibly cubic phases - and small unit cells are preferred for developing new structural materials. Figure 11-1 shows such crystal structures which are cubic (B2, D0 3 or L21? LI 2 , L/l 2 ) or nearly cubic, i.e. tetragonally distorted (Ll 0 , D0 22 ), and some of the examples - FeAl, NiAl, Fe 3 Al, Ni 3 Al, or TiAl - are candidate phases for developments of structural materials. However, even for such selected phases brittleness remains a problem and the causes of this are manifold (Liu et al., 1990; Baker and George, 1992). A material fractures in a brittle manner if there is no plastic deformation and stress relaxation at the crack tip, i.e. the yield stress is higher than the stress for cleavage or fracture. The reason for this may be an insufficient number and/or mobility of dislocations, and/or an insufficient number of slip systems. Since these parameters depend not only on crystal symmetry, but also on the specific characteristics of the particular phases, the aim of phase selection and composition variation by alloying must be dislocations with low energies and high mobilities and at least five independent deformation modes, i.e. dislocation slip systems and twinning systems, which are necessary for general homogeneous plastic deformation according to the Von Mises criterion (von Mises, 1928; Kocks, 1958, 1970; Kocks and Canova, 1981; Fleischer, 1988). Crystal anisotropy is an important factor in brittleness and criteria are used which rely on elastic moduli relations (Paxton and Pettifor, 1992). Brittle fracture may be a result of weak grain boundaries and other microstructure heterogeneities leading to localized deformation and stress concentrations. The disproportion of yield stress and fracture stress may also be due to a very low surface ener-

659

gy, which makes cleavage and cracking easy. This may be aggravated by the segregation of impurities which usually decrease the surface energies (Hondros, 1978). Impurities - in particular oxygen - may enter the material from the environment, and thus environmental embrittlement is an important issue for many intermetallics (Liu e t a l , 1990; Liu, 1991a). A detailed knowledge of the thermodynamics and kinetics of the particular alloy system is necessary for a reduction in the various embrittling effects. Intermetallics for any applications must be corrosion resistant in the respective environments. For high-temperature applications this means oxidation resistance in most cases. Oxidation resistance is provided by the presence of elements - in particular Cr, Al, or Si - which can form protective oxide layers (e.g., Aitken, 1967; Hindam and Whittle, 1982; Fitzer and Schlichting, 1983; Meier and Pettit, 1992). However, chromium oxide is volatile above 1000 °C and silicon oxide may form low-melting silicates. Thus aluminides are strongly favored for high-temperature applications. In order to act as protective layers, the scales must have sufficient density and adherence which may be improved by alloying with further elements (macroalloying) and, in particular, by adding minor amounts of "active" elements, e.g., Ti, Zr, or Hf (microalloying) (Rahmel and Schwenk, 1977; Hindam and Whittle, 1982). Furthermore, the scales must exhibit sufficient long-term stability which also means mechanical stability, i.e. they must tolerate strains, e.g. during creep, without being damaged (Riedel, 1982). In the case of insufficient oxidation resistance, protective coatings may be applied (see Pettit and Goward, 1983; Weatherill and Gill, 1988; Patnaik, 1989; Nicholls and Stephenson, 1991). However, the thermal stability of

660

11 Intermetallics

such coatings decreases with increasing temperature because of increasing reaction and diffusion rates, and therefore intermetallics for applications above 1100°C should be inherently oxidation-resistant. Finally, it must already be possible to prepare an intermetallic of sufficient quality in the laboratory, and from a practical point of view this is the basic requirement for any materials development. The processing of high-strength intermetallics is difficult because of their brittleness, and development of the necessary processing techniques is a very demanding task. It is, however, also a very important task, because a poor quality increases the apparent brittleness and reduces the strength which may then preclude any applications of the tested material.

11.3 Titanium Aluminides and Related Phases 11.3.1 Ti3Al 11.3.1.1 Basic Properties and Phase Diagram The titanium aluminide Ti3Al - often designated oc2 phase - crystallizes with the hexagonal ordered D0 1 9 structure (Ni3Sntype) which is shown in Fig. 11-9. The ratio of the lattice parameters c and a is c/a = 0.8 (Eckerlin etal., 1971; Villars and Calvert, 1991). The generally accepted value for the density is 4.2 g/cm3 (Fleischer, 1985; Munroe and Baker, 1988; Destefani, 1989), whereas for Ti3Al-base alloys the range 4.1-4.7 g/cm3 is given (see Table 11-2). This low density has made the titanium aluminides very attractive for materials developments. The characteristics of thermal expansion were described by Shashikala et al. (1989).

Table 11-2. Properties of alloys based on the titanium aluminides Ti3Al and TiAl compared with conventional titanium alloys and nickel-base superalloys (Morral, 1980; Lipsitt, 1985 a; Kim, 1989; Kim and Froes, 1990; Froes et al., 1991). Property Structure Density (g/cm3) Young's modulus at room temperature (GN/m2) Yield strength at room temperature (MN/m2) Tensile strength at room temperature (MN/m2) Temperature limit due to creep (°C) Temperature limit due to oxidation (°C) Tensile strain to fracture at room temperature (%) Tensile strain to fracture at high temperature (%) Fracture toughness Klc at room temperature (MN/m3/2)

Ti-base

Ti3Al-base

TiAl-base

Superalloys

A3/A2 4.5 95-115

D019/A2/B2 4.1-4.7 100-145

L1 0 /D0 19 3.7-3.9 160-180

A1/L12 7.9-9.1 195-220

380-1150

700-990

400-650

250-1310

480-1200

800-1140

450-800

620-1620

600 600 10-25

760 650 2-26

1000 900 1-4

1090 1090 3-50

12-50

10-20

10-60

8-125

high

13-42

10-20

25

661

11.3 Titanium Aluminides and Related Phases

1700

T D

»——•—

—' —. L

1

Sr 1S00 ( P Ti

• E 1300

N

1100CD Q.

E CD

900 (aTi) i /

700

/ /

500 o Ti

0

• Al

Figure 11-9. D 0 1 9 crystal structure of Ti 3 AL

(a)

Ti

/

j

' (a 2 ) /

s

—T Al,

• i

1

20

'.'TiAl

I

It*"

40

60

at.% Al

2

(A

80

100 Al

1400-

The temperature and composition dependence of the elastic constants of polycrystalline Ti3Al were studied (Schafrik, 1977), and, in particular, Young's modulus = 149 GPa, shear modulus = 58 GPa and Poisson's ratio = 0.29 were found for Ti3Al with 26at.% Al at room temperature. Young's moduli for Ti3Al-base alloys are in the range 100-145 GPa, which compares favorably with 96-110 GPa for conventional Ti-base alloys (Kimura et al., 1990; Froes et al., 1991). It has to be noted that in the case of Ti3Al (as in other cases) very different values have been measured for the Young's modulus, e.g. 77 MPa (Wolfenden et al., 1989). This emphasizes the fact that the elastic moduli depend sensitively on the alloy composition and microstructure including all artifacts which may be introduced during processing. Ti3Al has an extended composition range, as is visible in the commonly used Ti-Al phase diagram (Fig. 11-10 a), which, however, is still in discussion and has recently been revised, as shown in Fig. 11-lOb (Hellwig et al., 1992; Kainuma et al., 1994). It forms stable equilibria with the two disordered Ti phases, a-Ti with a hexagonal close-packed A3 structure and p-Ti with a b.c.c. A2 structure, and with the other important titanium alu-

13001200-

900 0

(b)

10 20 30 40 50 composition in at.% Al

Figure 11-10. (a) Ti-Al phase diagram according to Huang and Siemers (1989), Froes et al. (1991), Y. A, Chang et al. (1991), and Kim and Dimiduk (1991) and (b) revised version on the basis of new data according to Hellwig et al. (1992) and Kainuma et al. (1994).

minide TiAl, which is the subject of Sec. 11.3.2. The stabilities of the various phases in the Ti-Al system have been studied theoretically by first-principles calculations (Asta et al., 1993). It is noted that these equilibria, i.e. equilibrium compositions and transition temperatures, depend in a sensitive way on impurity contents, in particular on oxygen (Kahveci and Welsch, 1986; Huang and Siemers, 1989; Froes et al. 1991; Saunders and Chandrasekaran, 1992). Basic knowledge of the Ti-Al-O system is available (Glazova, 1965; Rah-

662

11 Intermetallics

mel and Spencer, 1991; Saunders and Chandrasekaran, 1992; Hoch and Lin, 1993), but the consequences of small oxygen additions on the Ti3Al equilibria have not yet been worked out in sufficient quantitative detail. Ti3Al has been alloyed with various substitutional and interstitial elements in order to control and optimize the mechanical properties and the corrosion behavior (Froes et al., 1991). With respect to the mechanical behavior, alloying with Nb, which substitutes for Ti, is very important (Rowe, 1990). The Ti-Al-Nb system has been analyzed recently (Hellwig, 1990; Kattner and Boettinger, 1992), and Fig. 11-11 shows a tentative version of the ternary Ti-Al-Nb phase diagram according to Hellwig (1990). It should be noted that the various equilibria in this diagram are not yet completely understood and are still in discussion because of conflicting experimental observations. In particular, additional phases, i.e. the orthorhombic O phase Ti 2 AlNb, oo structures, and a T 2 phase have been observed, and used, in the 1000°C

\

\

Ji)Al3

Nb

20

>Al 4 0

60

80

Al

at.% Al — Figure 11-11. Isothermal Ti-Al-Nb phase diagram section at 1000 °C where the p' phase with a B2 structure transforms into an co phase at lower temperatures (Hellwig, 1990).

Ti-rich corner besides the B2 phase (Rowe etal., 1991; Rowe, 1990; Koss etal., 1990; Perepezko, 1991a; Muraleedharan etal., 1992a, b; Bendersky etal., 1992; Hsiung and Wadley, 1992). The limited knowledge of the respective phase equilibria has been overviewed (Kim and Froes, 1990; Das et al., 1993 a). An improved understanding of these phase equilibria has been achieved by diffusion couple experiments and diffusion path analyses from which diffusion data can be obtained (Ma et al., 1992). The effects of alloying with Mo and V on the phase relationships have been studied recently (Das etal., 1993b; Ma and Dayananda, 1993). 11.3.1.2 Microstructure and Mechanical Behavior

Figure 11-12 shows the temperature dependence of the strength and plastic deformability of single-phase polycrystalline Ti3Al according to Lipsitt et al. (1980). It can be seen that this phase is brittle with practically no deformability at low temperatures up to 600 °C. Above this temperature plastic deformation was observed which was, however, still paralleled by intergranular cracking. Correspondingly, the fracture strength is nearly 600 MPa up to 600 °C. Above that temperature thermally activated softening occurs, which makes plastic deformation possible, and the resulting yielding leads to yield strengths below the fracture strength. The micromechanisms of deformation and in particular the dislocation reactions have been analyzed in detail and have been discussed with respect to strength and ductility (Koss et al., 1990; Kim and Froes, 1990; Yamaguchi and Umakoshi, 1990; Froes et al., 1991; Umakoshi et al., 1993 a, b). In the D0 1 9 structure five independent slip systems are possible (Kim and Froes,

11.3 Titanium Aluminides and Related Phases

663

1000

a E

200

400 temperature

600

800

Figure 11-12. Tensile strength and plastic deformability as a function of temperature for single-phase, polycrystalline, stoichiometric Ti3Al; = fracture stress (below 600 °C) or ultimate tensile strength, — • — = yield stress, = apparent deformability including microcracking, = estimated deformability without microcracking (Lipsitt etal., 1980).

1000

in °C

1990), which would satisfy the Von Mises criterion for uniform deformation (von Mises, 1928). However, single-crystal studies have shown that the yield stresses for the different slip systems are widely different (Minonishi, 1991; Nakano and Umakoshi, 1993), and thus not all five slip systems are activated during the deformation of polycrystalline Ti3Al. Only a little tensile ductility has been found for basal slip, whereas prism slip results in very large tensile elongations (Inui et al., 1993). Since there is no stress-relieving twinning, as in hexagonal metals, the insufficient number of slip systems, together with the observed planarity of slip, leads to strain incompatibilities and stress concentrations at grain boundaries from which cleavage fracture results (Koss et al., 1990; Kim and Froes, 1990; Froes etal., 1991). It is noted that one of the possible slip systems in single crystals shows an anomalous temperature dependence for the yield stress (Minonishi, 1991; Umakoshi etal., 1993 b), i.e. the yield stress increases with rising temperature until a maximum is reached. Such an anomalous temperature dependence is characteristic of various intermetallics and has been analyzed in much detail for the well-known case of Ni 3 Al, which is discussed in Sec. 11.4.1.2. The findings for Ni 3 Al, however, do not

apply in the case of Ti3Al because of different dislocation configurations (Minonishi, 1991). It has to be emphasized that the effect can only be observed in appropriately oriented, monocrystalline Ti3Al because the respective slip system is not activated in polycrystalline Ti3Al. The aim of the various materials developments based on Ti3Al is to improve both the strength and ductility by alloying with further elements and by controlling the microstructure (Rowe, 1990; Froes et al., 1991). The most effective element for improving ductility is Nb, as is illustrated in Fig. 11-13, and indeed Ti3Al-base alloys of engineering significance, i.e. oc2 alloys and super-oc2 alloys, contain 10-30 at.% Nb (Rowe, 1990; Froes etal., 1991). In spite of the extensive development work and the related numerous studies of microstructure and properties, the mechanism by which Nb improves the ductility of such Ti3Al-base alloys is not yet clear (Froes et al., 1991). Small amounts of Nb, which substitutes for Ti, lead to the activation of more slip systems which, however, has only a small effect on the ductility. Larger amounts of Nb result in the formation of further phases, i.e. (3-Ti in the disordered state with an A2 structure or in the ordered state with a B2 structure and/or the already mentioned orthorhombic O

664

11 Intermetallics

0

200

400

600

800 1000

femperat-ure in °C

Figure 11-13. Temperature dependence of strength for Ti3Al ( ) and Ti-24at.% Al-11 at.% Nb ( ), and of tensile elongation for Ti3Al ( ) and (Ti-24at.% Al-11 at.% Nb (-•-•-•-) (Rowe, 1990).

phase, which limit the slip length and have a significant, beneficial effect on the ductility. The dislocation configurations in the O phase have recently been studied (Douin etal., 1993). The alloys with good combinations of strength and ductility are multiphase intermetallic alloys with complex microstructures which may contain primary oc2 grains, fine secondary oc2 Widmanstatten platelets or laths, (3 grains in the disordered A2 or the ordered B2 state, co phase grains, and perhaps O phase grains (Koss et al., 1990; Rowe, 1990; Kim and Froes, 1990; J. M. Larsen etal., 1990; Froes etal., 1991). In any case, the mechanical behavior depends in a rather sensitive way on the distributions of the various phases, i.e. on the number, size, shape, composition, crystal structure, interface structure, and neighborhood relationships of the various grains. The phase distribution can be varied appreciably by proper selection of the

alloy composition and by thermomechanical treatments (Koss etal., 1990; Rowe, 1990; Kim and Froes, 1990; J. M. Larsen et al., 1990; Froes et al., 1991, 1992). Thus careful control of the processing is necessary for the optimization and reproducible production of oc2 alloys. It was found in particular that the optimum alloy microstructures are different for different mechanical properties, e.g. tensile strength, tensile ductility, creep resistance, fatigue crack growth resistance, and high cycle fatigue resistance (Rowe, 1990). The various toughening mechanisms have been reviewed recently with respect to crack initiation and crack growth in TiAl (Chan, 1993 a). Alloying with Nb improves most of the mechanical properties and the effect increases with increasing Nb. The only notable exception is the creep resistance which is a decisive property for high-temperature applications. Nb was found to reduce the creep resistance although it has also been found that the Nb-containing O phase may be beneficial for the creep resistance (Rowe, 1990; Kim and Froes, 1990; Froes et al., 1991; Nandy et al., 1993). Obviously the creep resistance is also a very sensitive function of the phase distribution, and current research aims at clarifying the conditions for optimum creep resistance (Hayes, 1991; Morris, 1991a; Huang and Kim, 1991; Es-Souni etal. 1991; Lupine etal., 1991; Onodera et al., 1991; Thompson and Pollock, 1991). The creep mechanisms have been overviewed (Yamaguchi and Umakoshi, 1990; Koss etal., 1990). The observed creep follows the known relationships for conventional disordered alloys, which has also been found for other intermetallics and will be discussed in Sec. 11.4.3.4 with NiAl alloys as examples. Diffusion data which are needed for analyzing and optimizing the creep behavior are

11.3 Titanium Aluminides and Related Phases

available, in particular for the Ti-Al-Nb system (van Loo and Rieck, 1973; Dayananda, 1992). Other alloying elements for improving the strength are Cr, Ta, and Mo. The latter is also advantageous for the creep resistance (Froes etal., 1991). Minor alloying additions of Fe, C and Si affect the creep behavior significantly, with Fe having the most deleterious effect (Rowe, 1990). Besides the cited elements, V and Sn were used for improving the properties. Alloying with Zr increases both the strength and the ductility, and microalloying with Y and B has been used to control the grain size and improve the ductility and workability (Froes etal., 1991). Very fine and stable grain sizes can be produced by rapid solidification of alloys with fine dispersions of rare earth oxides, e.g. Er 2 O 3 (Suryanarayana etal., 1991). In such alloys the small grain size improves the ductility, whereas the dispersoids enhance the strength at the expense of ductility. Similar effects can be produced by the precipitation of strengthening phases. This has been exemplified by materials developments which rely on the alloying of Ti3Al with Si to produce Ti 5 Si 3 as a strengthening second phase (see Wu et al., 1989; EsSouni et al., 1991, 1992a). Ti 5 Si 3 is a very hard and brittle phase with a complex hexagonal D8 8 structure, low density and a melting temperature above 2000 °C (see Sec. 11.11.3). Further alloying of such two-phase alloys with Nb is beneficial with respect to the mechanical behavior (Wagner et al., 1991; Es-Souni et al., 1991, 1992a). oc2 alloys of current engineering significance are Ti-24Al-llNb, Ti-25A110Nb-3V-lMo, Ti-25Al-17Nb-lMo, and Ti-23.5Al-24Nb (Froes etal., 1991), and the ranges of characteristic properties for these alloys are given in Table 11-2. The

665

mechanical behavior, i.e. strength, ductility, fatigue, fracture, and creep, has been discussed in detail in various recent reviews, in particular with respect to the various effects of the microstructure (Koss etal., 1990; Rowe, 1990; Kim and Froes, 1990; J. M. Larsen et al., 1990; Yamaguchi and Umakoshi, 1990; Froes etal., 1991; Kumpfert et al., 1992; Stoloff et al., 1993); see also the chapter by Umakoshi in Volume 6 of MST It is noted that superplastic deformation is also possible, during which dynamic grain growth with a decreasing volume fraction of Ti3Al occurs (H. S. Yang etal., 1992, 1993). 11.3.1.3 Environmental Effects

Exposure of Ti3Al or Ti3Al-base alloys to oxygen at higher temperatures leads to oxidation on the one hand and to oxygen dissolution in the alloy on the other (Rowe, 1990; Kim and Froes, 1990; Froes et al., 1991). Oxidation resistance would be expected if a protective A12O3, layer could be formed by selective oxidation. However, A12O3 is only slightly more stable than TiO and the activity of Ti is much higher than that of Al in Ti3Al (Meier and Pettit, 1992). Thus TiO is the stable oxide in contact with Ti3Al and further oxidation leads to the formation of TiO 2 , i.e. rutile. The thermodynamics of the system Ti-Al-O and the equilibrium conditions for the formation of the various oxide phases have been studied in detail (Rahmel and Spencer, 1991; Pajunen and Kivilahti, 1992; Li etal., 1992; Zhang etal., 1992b; Saunders and Chandrasekaran, 1992). According to these thermodynamic considerations, the observed oxidation behavior is complex leading to a layered oxide scale structure with TiO 2 on the outside and oxides with higher metal contents underneath, including A12O3 (Khobaib and

666

11 Intermetallics

Vahldiek, 1988; Shida and Anada, 1993). The kinetics follows a parabolic rate law and the rate constants of Ti3Al oxidation are only slightly smaller than those of typical TiO 2 formers, i.e. conventional Ti alloys (Choudhury etal., 1976; Perkins et al., 1989; Meier et al., 1989; Welsch and Kahveci, 1989). Nb increases the oxidation resistance to such an extent that the obtained rate constants are intermediate between those of TiO 2 formers and A12O3 formers (Choudhury et al., 1976; Khobaib and Vahldiek, 1988; Wiedeman et al. 1989; Welsch and Kahveci, 1989; Rowe, 1990; Kim and Froes, 1990). The oxidation resistance of Ti3Al-base alloys may also be improved by alloying with Mo or Ta (Froes etal., 1991; Schaeffer, 1993). The obtained oxidation resistances are not yet sufficient and limit the application of Ti3Al-base alloys at high temperatures, as is shown in Table 11-2. Therefore the coating of these alloys is studied extensively. In particular, pack-cementation techniques are optimized for aluminizing the alloy surface region to Al3Ti and avoiding surface cracking (Kung, 1990; Smialek etal., 1990 b). Besides scale formation at the surface, oxygen diffuses into Ti3Al as a solute since Ti3Al has a comparatively high solubility for oxygen. This leads to embrittlement, i.e. the strength is increased and the ductility is decreased, and to crack formation at the surface (Balsone, 1989; Rowe, 1990; Kim and Froes, 1990; McKee, 1993). The rate of embrittlement and crack formation is controlled by diffusion processes which can only be slowed down by protective coatings. It is noted that in spite of the inward diffusion of oxygen, internal oxidation with the formation of fine oxide particle dispersions has not been reported. The fact that classical internal oxidation does not occur, even though a protective oxide

layer has not formed or has broken down, is a common feature of many intermetallics and results from the particular thermodynamic and kinetic conditions in the respective alloy systems (Meier and Pettit, 1992). Another important environmental element is hydrogen (Chu and Thompson, 1991; Eliezer et al., 1991). It stabilizes the P phase in the Ti-Al system, i.e. it affects the thermodynamics of the system and thereby the microstructure evolution appreciably. It is indeed used as a temporary alloying element during thermochemical processing of conventional a/|3 titanium alloys for refining the microstructure. The solubility of H is highest in p-Ti, lower in oc-Ti and still lower in Ti3Al. Nevertheless, Ti3Al still dissolves significant amounts of H at high temperatures, which is precipitated as hydrides during cooling because of the lower solubility at lower temperatures (Chu and Thompson, 1991; Eliezer etal., 1991). Structure, shape and orientation relationships of the hydrides have been reported (Shih et al., 1989; Gao et al., 1990). Like other alloys, Ti3Al-base alloys are embrittled by both dissolved hydrogen and precipitated hydrides which leads to hydrogen-induced cracking (Eliezer et al., 1991; Gao etal., 1990; Chu and Thompson, 1991, 1992; Chu et al., 1992; Thompson, 1992, 1993). However, with special hydrogen loading conditions and heat treatments, i.e. thermochemical treatments, the strength, ductility and toughness can also be improved (Chu and Thompson, 1991). In particular, a hydrogen-tolerant microstructure has been developed for not too high hydrogen contents (Chan, 1992, 1993 b). More work is necessary here in order to understand the microstructural mechanisms and utilize thermochemical processing for the optimization of the mechanical behavior (Eliezer etal., 1991; Thompson, 1992).

11.3 Titanium Aluminides and Related Phases

11.3.1.4 Applications The developments of Ti3Al-based oc2 alloys have progressed in such a successful way that the a 2 alloys have been on the brink of commercialization for several years (Lipsitt, 1985 a, b, 1993; Fleischer e t a l , 1989a; Dimiduk etal., 1992; Froes et al., 1992). a 2 alloys are being produced in ingots of up to 4500 kg and are available as cast shapes or wrought forms including billets, bars, plates, and sheet (Peacock, 1989; Wittenauer etal., 1989; Chesnutt, 1990; Bassi et al., 1991). Likewise the powder metallurgy methods which have been developed for conventional titanium alloys are readily adapted, and in addition other powder metallurgy methods - in particular reactive processing - have been applied to Ti3Al-base alloys (Froes etal., 1991). a 2 alloys can be machined, sheet can be deformed superplastically and diffusion bonded to form complex sheet parts, and the various manufacturing processes can be accomplished on conventional equipment (Lipsitt, 1985a; Wittenauer etal., 1989; Bassi et al,, 1991). Joining studies have shown that ot2 alloys can be joined successfully by both diffusion bonding and linear friction welding, whereas with fusion welding problems of microstructure control may be encountered (Threadgill and Baeslack, 1991; Baeslack etal., 1988; Cieslak et al., 1990). Components for applications in flying gas turbines have been produced, they have been engine tested as static structures and they are nearing engine qualification (Lipsitt, 1985 a, b; Dimiduk et al., 1992). Examples are combustor swirlers, compressor casing sections, support rings, and afterburner nozzle seals. In spite of these successes the a 2 alloy components are not yet flying (Froes et al., 1991), i.e. they are not yet implemented

667

because expectations have risen with respect to strength and oxidation resistance, there are problems of ignition in some compressor applications, there is environmental embrittlement, and finally there are economic barriers (Dimiduk et al., 1992). Thus widespread use of oc2 alloys is expected only after further alloy and coating developments. Nevertheless, the a 2 alloys are being considered for many future aerospace applications (Roland, 1989; Dimiduk etal., 1991; Dauphin etal., 1991; Froes etal., 1991). Major initiatives are programs for developing piloted aerospace vehicles with hypersonic speeds, and examples are the U.S. National Aero-Space Plane (NASP) (Ronald, 1989) and the German Sanger Project (Kuczera et al., 1991; Bunk, 1992). Here specific strength, i.e. strength per unit density or specific gravity, is of primary interest and alloys to be selected have to compete with metallic and nonmetallic high-strength materials. In order to meet the design criteria, oc2 alloys are considered as matrix materials of intermetallic matrix composites, and respective materials developments are in progress (Bassi et al., 1991; Ronald, 1989; Bowman and Noebe, 1989; Destefani, 1989; Stephens, 1990; J. M. Larsen etal., 1990; MacKay etal., 1991; Norman etal., 1990; Feng and Michel, 1991; Jha et al., 1991 a, b; Kumar, 1991; Froes etal., 1991; Bryant etal., 1991; Marshall et al., 1992). Besides the problems of mechanical compatibility of the composite constituents, which are already present in conventional multiphase alloys, there is the problem of chemical stability which may be affected by reactions between the composite constituents (Norman etal., 1990). Thus composite developments aim at the optimization of composition and microstructure with respect to composite sta-

668

11 Intermetallics

bility, tensile properties, thermal and mechanical fatigue resistance, creep resistance, and corrosion resistance, as is exemplified by MacKay et al. (1991). Finally, the success of such developments depends on economics, i.e. total life cycle costs. 11.3.2 TiAl 11.3.2.1 Basic Properties and Phase Diagram The titanium aluminide TiAl - often designated as y phase - crystallizes with the tetragonal L l 0 structure (CuAu-type) which is shown in Fig. 11-1. The LI 0 structure results from ordering in the f.c.c. lattice (Al), i.e. it is basically a cubic structure which is tetragonally distorted because of the particular stacking of the atom planes, as is seen in Fig. 11-1. The ratio of the lattice parameters c and a is c/a = 1.015 at the stoichiometric composition and the density is 3.76 g/cm3 (Kim and Dimiduk, 1991), whereas for TiAl-base alloys the range 3.7-3.9 g/cm3 is given (see Table 11-2). This density is still lower than that of Ti3Al and has made the titanium aluminides most attractive for materials developments. The temperature dependence of the elastic constants of polycrystalline TiAl was studied (Schafrik, 1977), and Young's modulus = 174 GPa, shear modulus = 70 GPa, and Poisson's ratio = 0.23 were found for stoichiometric TiAl at room temperature, i.e. the elastic constants are larger and Poisson's ratio is smaller than for Ti 3 Al. Elastic constants and fault energies of TiAl have been studied theoretically by first-principles total-energy investigations and have been discussed with respect to mechanical behavior (Yoo e t a l , 1991; Lee and Yoo, 1990; Fu and Yoo, 1990; Yoo, 1991; Woodward etal., 1993; Yoo and Fu, 1991, 1993). Young's moduli for

TiAl-base alloys are in the range 160-180 GPa which is only 10-20% lower than that of the superalloys (see Table 11-2). Recently, it has been found by ab initio calculations that deviations from stoichiometry are due to accommodated antistructure atoms, i.e. constitutional disorder, instead of vacancies in the sublattices, and that the concentration of thermal vacancies is comparatively low because of the high formation energy (Fu and Yoo, 1993). The self-diffusion of Ti in TiAl has been studied (Kroll et al., 1992). TiAl has an extended composition range and is stable up to the melting point, as is shown in the Ti-Al phase diagram (Fig. 11-10). The variation in Al content between the solubility limits leads to constitutional disorder with excess Ti or Al atoms on Al or Ti sites, respectively, i.e. antistructure atoms on antisites without constitutional vacancies, and it results in a corresponding variation in the c/a ratio, i.e. the tetragonality, between 1.03 for maximum Al content and 1.01 for minimum Al content (Yamaguchi and Umakoshi, 1990; Kim and Dimiduk, 1991; Shirai and Yamaguchi, 1992). The stability of the various phases in the Ti-Al system has been studied theoretically by first-principles calculations (Asta et al., 1993). It is noted that the Ti-Al phase diagram has been in discussion for many years and presently the version in Fig. 11-10 a is generally used on the basis of various recent studies (Y A. Chang et al., 1991; Shull and Cline, 1990; Perepezko, 1991a, b; Huang and Siemers, 1989; Kim and Dimiduk, 1991; Rowe und Huang, 1988; Kattner etal., 1992; Jones and Kaufman, 1993; Anderson et al., 1993). Nevertheless, the phase equilibria with TiAl at high temperatures are still not clear, and a new, careful study has resulted in a revised diagram, as shown in Fig. ll-10b (Hellwig et al., 1992;

11.3 Titanium Aluminides and Related Phases

Kainuma etal., 1994). As already noted with respect to Ti3Al, the phase equilibria, i.e. equilibrium compositions and transition temperatures, depend in a sensitive way on impurity contents, in particular on oxygen (Kahveci and Welsch, 1986; Huang and Siemers, 1989; Froes e t a l , 1991; Saunders and Chandrasekaran, 1992). TiAl has been alloyed with various substitutional and interstitial elements in order to control and optimize the mechanical properties and the corrosion behavior (Hashimoto etal., 1986a, b, 1991; Kasahara et al., 1987; Tsujimoto and Hashimoto, 1989; Froes etal., 1991; Kim and Dimiduk, 1991). Alloying elements have different effects on the extension of the TiAl field in the respective ternary phase diagrams, as is shown schematically in Fig. 11-14. In this crude approximation, V, Mn, and Cr apparently act as Al substituents or as both Ti and Al substituents whereas Nb, Ta, Zr, Mo, and W act as Ti substituents (Kim and Dimiduk, 1991; X. F.Chen etal., 1992). The site preference of such alloying additions in the TiAl lattice and their effect on the tetragonality have been studied recently by ab initio calculations and the results are in partial agreement with the above experimental findings (Erschbaumer et al., 1993). In detail, the effects are complex and also depend on the amount of alloying addition, as is revealed by the available ternary phase diagrams for TiAl with Nb (see Fig. 11-11), Mn, Zr, V, or Ag (Hashimoto et al., 1986a, b; Kasahara et al., 1987; Tsujimoto and Hashimoto, 1989; Kim, 1989; Ahmed and Flower, 1992). As already mentioned with respect to Ti3Al (Sec. 11.3.1.1), the phase equilibria in the Ti-Albase systems are not yet fully understood and the character of additional phases is the subject of ongoing research (Jackson and Lee, 1992; Jackson, 1993; Nakamura

20

30

40

50

669

60

Al

Al content" in at % Figure 11-14. Schematic isothermal section at 900 °C of the ternary phase diagram Ti-Al-M with TiAl lobe I for M = Nb, Ta, Mo, or W and lobe II or III for M = V, Cr, or Mn; area C marks the composition range of commercial interest (Kim and Dimiduk, 1991).

etal., 1993). An improved understanding of these phase equilibria has been achieved by diffusion couple experiments and diffusion path analyses from which diffusion data can be obtained (Ma etal., 1992). The effects of alloying with Mo and V on the phase relationships have recently been studied (Das etal., 1993b; Ma and Dayananda, 1993). 11.3.2.2 Microstructure and Mechanical Behavior

The variation of strength and ductility with temperature is similar to that of the Ti-rich Ti3Al in Fig. 11-12. Figure 11-15 shows the temperature dependence of the strength and plastic deformability of single-phase polycrystalline TiAl according to Lipsitt et al. (1975). It can be seen that this phase is brittle with practically no deformability at temperatures up to 700 °C, and only above that temperature was plastic deformation observed. Correspondingly, it has a fracture strength of nearly 500 MPa up to about 700 °C. Above that temperature, thermally activated softening occurs making plastic deformation possible, and the resultant yielding leads to yield strengths below the fracture strengths.

670

11 Intermetallics

0

200

400

600

800

1000

temperature in °C

Figure 11-15. Tensile strength, i.e., fracture stress, ultimate tensile strength (UTS), 0.2% yield stress, and plastic deformability, i.e., tensile elongation, as a function of temperature for single-phase, polycrystalline TiAl with 54 at.% Al (Lipsitt et al., 1975).

The micromechanisms of deformation, and in particular the dislocation reactions, have been analyzed in much detail and have been discussed with respect to strength and ductility (Shechtman et al., 1974; Lipsitt etal., 1975; Yamaguchi and Umakoshi, 1990; Greenberg etal., 1991; Kim and Dimiduk, 1991; Yoo and Fu, 1991; Simmons etal., 1993; Denquin and Naka, 1993; Li and Whang, 1993; Whang and Hahn, 1993); see also the chapter by Umakoshi in Volume 6 of MST. Plastic deformation is achieved by the movement of single dislocations and superdislocations which both exhibit sessile components. In addition, twinning is an important deformation mechanism. Low dislocation mobility is the main cause for the brittle fracture behavior at lower temperatures since it hinders the formation of a plastic zone at the crack tip, which relaxes the stress concentrations (Yoo and Fu, 1991). A recent study has been directed at the effect of grain size on the mechanical behavior, which cannot be described by the Hall-Petch relation (Imayev et al., 1993 a, b).

Single-crystalline TiAl shows an anomalous, positive temperature dependence for the yield stress, i.e. the yield stress increases with rising temperature until a maximum is reached at about 600 °C, and only above this peak temperature is there the normal negative temperature dependence (Kawabata et al., 1985; Huang and Hall, 1991 a; Y G. Zhang et al., 1991; Veyssiere, 1991; Stucke et al., 1993). The anomaly results from thermally activated cross-slip of glissile superdislocations with comparatively high mobility to slip planes where not only the energy of the superdislocation is lower, but also the mobility because of a high Peierls stress, i.e. the superdislocations are immobilized by thermally activated cross-slip pinning (Kawabata et al., 1991a; Fu and Yoo, 1990; Yoo and Fu, 1991). The energy difference between the dislocation configurations, i.e. the driving force for cross-slip, originates from the orientation dependence of the antiphase boundary (APB) energy and from the torque force between interacting dislocation partials due to the anisotropy of elasticity. These effects, i.e. high and anisotropic APB energies, strong anisotropy of elasticity, and high Peierls stresses opposing dislocation motion, are related to the strong directional bonding between Ti and Al atoms (Fu and Yoo, 1990; Yoo and Fu, 1991). This anomalous temperature dependence of the yield stress is not always observed in polycrystalline TiAl alloys, as shown in Fig. 11-15. The reason for this is that there are additional strengthening effects at lower temperatures - in particular grain boundary hardening - which increase the strength to such an extent that the strength peak is no longer discernible (Huang and Hall, 1991 a). A recent study is directed at the effects of grain size, Al content, and impurity content on the flow

671

11.3 Titanium Aluminides and Related Phases

stress anomaly in TiAl (Sriram et al., 1993). Similar anomalies have also been observed for other intermetallics, and the best-known example is Ni 3 Al, which is discussed in Sec. 11.4.1.2. The mechanisms which cause the anomalous temperature dependence in Ni3Al and in TiAl have similar characteristics, however, the particularities are different because of the different crystal structures and dislocation configurations (Kawabata etal., 1991a). The major problem in the use of TiAl is its low ductility at room temperature. The low ductility of single-phase TiAl (with more than 50 at.% Al) is not improved by alloying additions of, e.g. V, Nb, Cr, W, or Mn (J. M. Larsen et al., 1990; Kim and Dimiduk, 1991). A reduction in the Al content improves the ductility. If the Al content is reduced below 50 at.% to form Ti3Al as a second phase besides the TiAl, the resulting intermetallic two-phase alloys show ductilities of a few percent, which are acceptable for applications (Huang and Hall, 1991b; Kim and Dimiduk, 1991; Froes etal., 1991; Nonaka etal., 1992). This is illustrated in Fig. 11-16 which shows the hardness and ductility at room temperature of single-phase and two-phase TiAl alloys. According to these data, 48 at.% Al is the optimum composition for maximum room-temperature ductility of TiAl alloys. The different hardness curves for the as-cast and annealed conditions indicate the strong effect of microstructure variations on the mechanical behavior. The aim of the various materials' developments based on TiAl is to improve both the strength and ductility, first by controlling the microstructure and second by alloying with further elements. The microstructure, i.e. the distribution of phases, can be varied within wide limits by appropriate heat treatments and thermomechanical processing (Koeppe et al.,

45

BO

SS

60

Al content in at.%

Figure 11-16. Vickers hardness and tensile elongation at room temperature as a function of Al content for single-phase and two-phase TiAl alloys (Kim, 1989).

1993). During processing, dynamic recrystallization occurs which has been studied systematically (Lee etal., 1993). Depending on the particular heat treatment, coarse equi-axed TiAl grains with stringers of fine TiAl grains pinned by Ti3Al particles, or a fine duplex microstructure consisting of TiAl grains and lamellar grains, or a nearly lamellar microstructure with coarse lamellar grains and minor amounts of fine TiAl grains, or a fully lamellar microstructure with large lamellar grains can be formed. Lamellar grains in the various microstructures are composed of alternating plates of TiAl and Ti3Al (Kim and Dimiduk, 1991). The various microstructures result in rather different mechanical properties and offer good possibilities for optimization (Y.-W. Kim, 1992), as is shown in Figs. 11-17 and 11-18. It is noted that the small grain size of the duplex alloys is advantageous for short-term strength and ductility, but detrimental to creep resistance, i.e. long-term strength, because of grain boundary sliding. Such thermomechanically processed, fine-grained alloys can be

672

11 Intermetallics 30

20-

10-

C

reciprocaTcree'p rate

I

n

i

IV

microstructure Figure 11-17. Fracture toughness, tensile strength, and tensile elongation at room temperature, and reciprocal secondary creep rate measuring the creep resistance as a function of microstructure (schematic) for two-phase TiAl alloys (Kim and Dimiduk, 1991).

200 800-

100 700-

50 600-

20

S00-

10

400-

-5

300-

a en c

-2

200-

1

100BDTT

0

0

200

400 600

800 1000

0.3

temperature in°C Figure 11-18. Temperature dependence of ultimate tensile strength (UTS), tensile yield strength (YS), brittle-to-ductile transition temperature (BDTT), and tensile elongation (El) for two-phase TiAl alloys with various microstructures produced under various processing conditions, in particular thermomechanical processing (TMP) (Kim and Dimiduk, 1991).

deformed superplastically (Cheng et al., 1992). The character and orientation of the phase boundaries is of particular importance for the deformation behavior, which is controlled by slip and twinning (Wunderlich et al., 1993; Seeger and Mecking, 1993; Appel et al., 1993). The creep behavior of single-phase and two-phase TiAl alloys, which includes twinning, has been studied in detail (Feng etal., 1990b; Huang and Kim, 1991; Oikawa, 1992; Maruyama et al., 1992; Hayes and London, 1992; Jin and Bieler, 1993; Bartels etal., 1993). Analysis of the observed creep with respect to the contributions of the constituent phases has been the subject of Bartholomeusz et al. (1993). Only a few studies have been directed at the fatigue behavior of TiAl alloys (Feng et al.; 1990 a; Aswath and Suresh, 1991; Dowling et al., 1991; Froes etal., 1992; Rao etal., 1992; Soboyejo etal., 1993; Stoloff et al., 1993). It has to be noted that particular microstructures may have contrasting effects on the ductility and toughness, as is demonstrated by the properties of the lamellar microstructures in Fig. 11-17. The various toughening mechanisms have been reviewed recently with respect to crack initiation and crack growth in TiAl (Chan, 1993 a). Two-phase TiAl alloys with a very special microstructure and striking mechanical behavior have been produced by using a special crystal-growth technique (Yamaguchi, 1991; Umakoshi et al., 1992; Inui et al., 1992; Umakoshi and Nakano, 1993; Yamaguchi and Inui, 1993). For example, the alloy with 49.3 at.% Al consists of one or two lamellar grains which are composed of lamellae of the major constituent TiAl and of the secondary phase Ti 3 Al, and the TiAl lamellae contain large numbers of thin twins. Accordingly, alloys with this microstructure are called polysynthetically

11.3 Titanium Aluminides and Related Phases

twinned (PST) crystals. This special microstructure gives rise to a mechanical behavior which depends sensitively on the orientation of the crystal with respect to the loading direction. There is a hard deformation mode with shearing across the lamella boundaries and an easy mode with shearing parallel to the lamella boundaries. The latter allows tensile elongations of 20 % at room temperature. Besides carefully controlled heat treatments and thermomechanical processing, small additions of alloying elements - usually 1-3 at.% - are used for the optimization of the mechanical behavior of the twophase TiAl alloys (Kim, 1989; Tsujimoto and Hashimoto, 1989; Yamaguchi and Umakoshi, 1990; Hashimoto et al., 1991; Froes etal., 1991, 1992; Kawabata etal., 1991b, 1993; Kim and Dimiduk, 1991; Huang etal., 1991a; Strangwood etal., 1992; Tsujimoto et al., 1992; Huang, 1993; Hashimoto and Kimura, 1993). V, Hf, Cr, and Mn increase the ductility significantly and produce solid-solution strengthening, with Cr being most effective and Mn being least effective. Nb, Ta, and W also produce solid-solution strengthening, but they decrease the ductility. Interstitial elements such as C and N affect the ductility, depending on the Al content and pre-treatments, and in particular they improve the creep resistance (Kawabata etal., 1991b; Kim and Dimiduk, 1991). The effects of the various elements have been attributed to changes in the electronic distribution and bonding character leading to changes in the lattice unit cell dimensions, tetragonality and site occupancy. However, these elements also affect the microstructures because of their effects on the thermodynamics of the Ti-Al system, and it is difficult to separate the various effects. Thus a clear understanding of the specific roles of the various alloying elements has not yet

673

been established (Kim and Dimiduk, 1991). A comparatively simple mechanism is precipitate strengthening which has been studied by Nemoto et al. (1992). The various materials developments in progress have led to a spectrum of multinary alloys which differ in Al content, amount of alloying elements, and processing (Tsujimoto and Hashimoto, 1989; Hashimoto et al., 1991; Froes etal., 1991; Kawabata etal., 1991b; Kim and Dimiduk, 1991; Tsujimoto etal., 1992). The obtained property ranges are illustrated in Table 11-2. A special Chinese development relies on alloying TiAl with Nb in order to obtain strong and oxidation-resistant, lightweight alloys (G. Chen et al., 1991,1992, 1993; W. Zhang et al., 1993). Alloying with Nb leads to atomic ordering in the L l 0 structure of TiAl, i.e. a new phase is produced, the crystal structure of which is not yet quite clear. It has already been emphasized in the preceding section that the phase equilibria of the titanium aluminides are very sensitive to interstitial impurities, in particular oxygen. It was shown recently that an almost single-phase TiAl alloy with 50 at.% Al can be obtained by using high-purity Ti and Al, whereas the conventional Ti-50 at.% Al alloy of low purity is two-phase with a significant amount of Ti3Al as well as TiAl (Murata et al., 1992, 1993). The high-purity TiAl is ductile with a fracture strain of more than 3 % in contrast to the low-purity TiAl with only about 1 % fracture strain. 11.3.2.3 Environmental Effects

The oxidation resistance of TiAl is higher than that of Ti3Al because of its higher Al content, but it is still orders of magnitude lower than that of typical A12O3 formers, e.g. NiAl (Choudhury

674

11 Intermetallics

etal., 1976). The oxidation resistance of TiAl relies on the formation of a protective A12O3 layer which is, however, only slightly more stable than TiO, as has already been remarked upon with respect to Ti3Al in Sec. 11.3.1.3 (Meier and Pettit, 1992). Because of the high Ti activity in Ti-rich alloys, TiO is the stable oxide in contact with the TiAl alloys with Al contents below 50 at.%, which is the case for alloys with maximum ductility (Kim and Dimiduk, 1991). Further oxidation leads to the formation of TiO 2 , i.e. rutile, which is not protective, and thus oxidation is a problem for TiAl alloys which have been optimized with respect to mechanical behavior. The thermodynamics and the equilibrium conditions for the formation of the various oxide phases have been studied in detail (Rahmel and Spencer, 1991; Zhang etal., 1992b). The oxidation behavior of such TiAl alloys is significantly improved by alloying with Nb, which results in the formation of a protective A12O3 layer with an improved resistance to spalling (Becker etal., 1993). Oxidation of TiAl alloys with sufficiently high Al contents, i.e. single-phase alloys with at least 50 at.% Al, leads to the formation of A12O3 scales with correspondingly low oxidation rates only at temperatures below about 1000 °C, whereas at higher temperatures complex scales develop with an outer rutile layer over a mixed layer of rutile and A12O3, with markedly increased oxidation rates (Meier and Pettit, 1992; Taniguchi etal., 1991a). Oxidation at the surface reduces the Al content of the TiAl to such an extent that a layer of brittle Ti3Al - usually with cracks - is formed. Below these scales and surface layers internal oxidation is observed, resulting in oxide dispersions in the bulk. The transition from low-temperature oxidation to high-temperature oxidation occurs in a

very narrow temperature range, and the reasons for this behavior are not yet understood (Meier and Pettit, 1992). The micromechanisms and kinetics of scale formation have been studied recently (Becker et al., 1992; Shida and Anada, 1993). The oxidation resistance of TiAl alloys can be improved by special pre-oxidation treatments (Suzuki etal., 1991). Alternatively, it can be improved by alloying with Nb, Ta, and W which, however, reduces the ductility; whereas V, Cr, and Mn, which are used for increasing ductility, reduce the oxidation resistance (Kim, 1989). TiAl-based alloys with high contents of Nb have been studied recently with respect to the conditions for protective scale formation, and indeed protective oxidation has been found at 1400 °C for an Al content of 50 at.% (Brady et al., 1993). In view of the oxidation problems of the titanium aluminides a coating has been proposed for providing sufficient oxidation protection, and various approaches have been studied (Nishiyama et al., 1990; Taniguchi et al., 1991 b; Yoshihara et al., 1991; Wu and Lin, 1993). Besides scale formation, oxygen embrittles TiAl alloys as do other interstitial impurities, i.e. N, C, and B (Kim and Dimiduk, 1991). It is noted that the solubility of oxygen in TiAl is lower than that in Ti3Al which, however, is present in most TiAl alloys and thus getters the oxygen (Yamaguchi and Umakoshi, 1990; Kim and Dimiduk, 1991). The yield stress is increased at low and high temperatures by both dissolved oxygen and precipitated A12O3 particles (Kawabata etal., 1992). Environmental embrittlement has been found for TiAl as for other intermetallics (Liu and Kim, 1992). This embrittlement is obviously caused by the presence of moisture in the test atmosphere since the ductility of an advanced TiAl alloy (with addi-

11.3 Titanium Aluminides and Related Phases

tions of Cr, Mn, and Si) has been found to be lowest in air and highest in pure oxygen. Such embrittlement effects depend sensitively on the temperature, strain rate, microstructure and phase distribution (Chan and Kim, 1993; Kim and Dimiduk, 1993). The environmental embrittlement of polysynthetically twinned TiAl (see the preceding section) is reduced by alloying with Cr, Mo or Mn (Oh et al, 1993a, b). The solubility of hydrogen in TiAl is still sufficiently high for hydride formation on cooling, which had not been detected in earlier reports (Thompson, 1992). No clear effects of hydrogen on the mechanical behavior have been found for single-phase TiAl, whereas two-phase alloys with Ti3Al showed embrittlement probably due to the gettering effect of Ti3Al. More work is necessary for a clear understanding of the embrittling mechanisms (Eliezer et al., 1991; Thompson, 1992). 11.3.2.4 Applications The TiAl alloy developments are less advanced than the Ti3Al-based a 2 alloy developments because of the more severe ductility problems. However, the TiAl alloys are regarded as more promising primarily because of their lower density and the higher application temperatures, and development activities are now concentrated on TiAl in various countries, e.g. in the U.S.A. (Kim and Dimiduk, 1991; Froes e t a l , 1991, 1992; Dimiduk et al., 1991; Huang et al., 1991 b), in Japan (Nishiyama etal., 1990; Yamaguchi, 1992; Matsuo, 1991; Nakao etal., 1991; Mabuchi etal., 1991; Kusaka, 1991; Yoshihara etal., 1991; Ogishi et al., 1991; Kawabata et al., 1991b; Fujitsuna etal., 1991; Tokizane et al., 1991; Hashimoto et al., 1991; Masahashi etal., 1991; Matsuo etal., 1991; Nakagawa etal., 1992; Tsujimoto etal.,

675

1992), Great Britain (Peacock, 1989), Russia (Bondarev etal., 1991; Imayev etal., 1993 b) and Germany (Sauthoff, 1990 a; Dogan etal., 1991; Dahms etal., 1991; Kuczera etal., 1991; Frommeyer etal., 1992). Processing is more difficult than for Ti3Al alloys because of the higher brittleto-ductile transition temperature, and in particular because the hot working temperatures are outside the range of conventional titanium processing equipment, which has limited the practical ingot size to about 200 kg (Peacock, 1989; Chesnutt, 1990; Froes etal., 1991; Kusaka, 1991; Matsuo, 1991). Present efforts are directed at improving the alloy quality by improved melting processing and thermomechanical treatments (Sakamoto etal., 1992; Lombard etal., 1992; Szaruga etal., 1992; Mouldckhues and Sahm, 1992; Guan etal., 1994; Austin and Kelly, 1993; London et al., 1993; T. P. Johnson et al., 1993; Takeyama etal., 1993). The problem of upscaling heat treatments for large production billets has been considered (Semiatin et al., 1993). Besides casting and hot working, powder metallurgy has been used for component fabrication, though there are problems with quality, availability, and the cost of powder (Froes et al., 1991; Chesnutt, 1990; Dahms etal., 1991, 1993; Tokizane etal., 1991; Fuchs, 1993; Oehring et al., 1993). Cast TiAl alloys can be forged isothermally (Fujitsuna etal., 1991; Clemens et al., 1993 b), fine grain powder metallurgy material can be formed superplastically (Tokizane etal., 1991; Matsuo, 1991) as well as special cast and thermomechanically treated material (Tsujimoto et al., 1992), sheet of 1-2 mm thickness can be produced by twin-roll casting, special processes have been developed for rolling, and machining is possible as well as joining by diffusion bonding and welding (Matsuo,

676

11 Intermetallics

1991; Kim and Dimiduk, 1991; Fujitsuna etal., 1991; Tokizane etal., 1991; Wurzwallner etal., 1993). The diffusion bonding process has been studied in particular detail (Yan and Wallach, 1993). With respect to turbine engine applications, component fabrication techniques have been demonstrated for airfoils and compressor cases (Lipsitt, 1985 a; Chesnutt, 1990; Yamaguchi and Umakoshi, 1990; Kim and Dimiduk, 1991; Bondarev et al., 1991). However, much more development work is necessary to overcome the problems of ductility and toughness, corrosion resistance, strength and costs. Nevertheless, TiAl alloys are being considered for future aerospace applications (Dimiduk et al., 1991; Kim and Dimiduk; 1991; Matsuo, 1991), and in particular for applications in the U.S. National AeroSpace Plane (NASP) (Ronald, 1989) and the German Sanger Project (Kuczera et al., 1991; Bunk, 1992), which were mentioned earlier with respect to Ti3Al alloys in Sec. 11.3.1.4. Here specific strength, i.e. strength per unit density or specific gravity, is of primary interest and alloys to be selected have to compete with metallic and nonmetallic high-strength materials. In order to meet the design criteria, TiAl alloys are considered as matrix materials of intermetallic matrix composites, and respective materials developments are in progress (Anton, 1988; Christodoulou et al., 1988; Feng etal., 1990b; Norman etal., 1990; Rosier etal., 1990; Westwood, 1990; Bryant etal., 1991; Froes etal., 1991, 1992; Kumar, 1991; Kumar and Whittenberger, 1991, 1992; Mabuchi etal., 1991; Soboyejo etal., 1993; Sadananda and Feng, 1993). Here the additional problems of chemical and mechanical compatibility have to be solved, as was mentioned earlier with respect to Ti3Al alloys in Sec. 11.3.1.4. Deterioration of the tensile

strength of a TiAl/SiC composite by interface reactions has recently been studied (Ochiai et al., 1994). The kinetics of the phase reactions is determined by interdiffusion, which has been studied in the case of TiAl-Mo (Zhang etal., 1992a). It has been shown that both the creep resistance and the toughness of a TiAl matrix composite can be improved by using coated fibers with weak fiber/matrix interfaces (Weber et al., 1993). In any case, ductile reinforcing inclusions in composites offer the valuable possibility of improving the fracture toughness of the composite. The toughening effect of ductile reinforcements in brittle matrices has been studied both experimentally and theoretically for ceramics, e.g. A12O3 reinforced with Al, and metals (Sigl et al., 1988; Flinn etal., 1989; Ashby etal., 1989). Such a toughening effect has also been found for a TiAl/Nb model composite (Cao etal., 1989). Besides aerospace applications, interest is focused on automotive engine applications, e.g. valves and turbocharger rotors, which may be nearer to realization (Kim and Dimiduk, 1991). A successful development has led to a cast TiAl turbocharger rotor of about 4 cm diameter which compares favorably with both a superalloy Inconel 713C rotor and a ceramic SiN rotor with respect to performance under service conditions, and which is ready for use in high-performance passenger cars (Nishiy a m a e t a l , 1990). 11.3.3 Al3Ti and Other D0 22 Phases 11.3.3.1 Basic Properties and Phase Diagram The titanium trialuminide Al3Ti crystallizes with the tetragonal D0 22 structure which is shown in Fig. 11-1, and which is common to some other trialuminides, i.e.

11.3 Titanium Aluminides and Related Phases

A13M with M = V, Nb, Ta, and some other phases, e.g. Ni3V (Bauer, 1939; Villars and Calvert, 1991). The D0 22 structure is derived from the close-packed cubic LI 2 structure by the stacking of L l 2 cubes with periodic antiphase boundaries in between, and thus the D0 2 2 structure may be regarded as a tetragonally distorted, longperiod ordered, cubic structure (Bauer, 1939; Yamaguchi and Umakoshi, 1990). The D0 22 structure is still sufficiently simple for quantum-mechanical, first-principles calculations and indeed such ab initio calculations have been done for Al3Ti in order to study the bonding type and phase stability, and to obtain theoretical values for the elastic constants, fault energies and cleavage energies (Carlsson, 1991; Lin e t a l , 1991; Yoo and Fu, 1991). The experimentally determined elastic stiffness values are larger than the theoretical ones, and in particular the resulting experimental value of 216 GPa for the Young's modulus of polycrystalline Al3Ti at room temperature clearly surpasses those of the other titanium aluminides, and is of the order of that of the superalloys (see Table 11-2) (Nakamura, 1991; Nakamura and Kimura, 1991). Because of the high Al content of Al3Ti its density of 3.3 g/cm3 is still lower than that of the other titanium aluminides (Yamaguchi and Umakoshi, 1990). Furthermore, its oxidation resistance is significantly higher than that of Ti3Al and TiAl (Umakoshi etal., 1989; Subrahmanyam and Annapurna, 1986), and thus Al3Ti is a candidate phase for lightweight structural applications. It is also being considered as a coating material and its oxidation kinetics have been studied in detail (Smialek and Humphrey, 1992). Al3Ti melts incongruently at about 1340°C and is a line compound with a fixed composition according to the avail-

677

able phase diagrams (Glazova, 1965; Murray, 1988; Rowe and Huang, 1988; Huang and Siemers, 1989; Schuster and Ipser, 1990; Shull and Cline, 1990; Y. A. Chang etal., 1991; Perepezko, 1991a). It should be noted that the high-temperature equilibria are still in discussion. Al3Ti can be alloyed with the other D0 2 2 trialuminides to form ternary D0 22 phases which are no longer line compounds, but have a range of homogeneity, i.e. the composition can vary between the solubility limits which are given in the respective ternary phase diagrams (Sridharan and Nowotny, 1983; Hashimoto etal., 1986b; Kumar, 1990; Hellwig, 1990; Paruchuri and Massalski 1991; Weaver etal., 1991; Perepezko, 1991a). 11.3.3.2 Microstructure and Mechanical Behavior

Al3Ti The deformation behavior of Al3Ti has been studied and analyzed in detail (Yamaguchi et al., 1987, 1990; Wheeler et al., 1990; Morris and Lerf, 1991; Yamaguchi and Umakoshi, 1990; Shimokawa etal., 1991; and the chapter by Umakoshi in Volume 6 of MST). Figure 11-19 shows the temperature dependence of the yield stress for polycrystalline material with a grain size of about 1 mm, which was produced by vacuum induction melting and contains a small volume fraction of second phase particles (mostly Al) due to the line compound character of this phase. The strength at room temperature is low compared with the other titanium aluminides (see Table 11-2) at about 500 °C. The major deformation mode is twinning, which does not affect the D0 22 symmetry. The reason for the preponderance of twinning is the low mobility of dislocations, according to a theoretical study of the deformation be-

678 1

11 Intermetallics JW

1000-

Al3(Nb0J5Ti0J5)

k

£

A

Al3Nb\

500At 3 T i

\

0200 400

600

800 1000 1200 1400

remperarure in °C

Figure 11-19. Yield strength of Al3Ti under vacuum (Yamaguchi et al., 1987) and 0.2% proof stress of Al3Nb and Al3(Nb0 75 Ti 0 25) in air (Sauthoff, 1990b; Reip, 1991; Reip and Sauthoff, 1993) as a function of temperature (coarse-grained, polycrystalline alloys with a D0 22 structure, tested in compression at a strain rate of 10 4 s *).

havior of D0 22 crystals (Khantha et al., 1992). In spite of the observed microscopic plasticity, there is nearly no ductility at temperatures below 620 °C and cracks are nucleated readily at structure heterogeneities, i.e. inclusions, second phase particles, intersections of twins, and grain boundaries. The ductility of Al3Ti is improved by microalloying, i.e. by small additions of Zr and Hf as well as B and Li, and a fracture strain in compression of a few percent is obtained at room temperature (Yamaguchi etal., 1988). The effect of Zr and Hf is supposed to be due to a lowering of the stacking fault energy, since the phases Al3Zr and Al 3 Hf crystallize with a D0 2 3 structure which differs from the D0 2 2 structure by a longer stacking period of LI 2 cubes. Al3Ti is being considered for use as an oxidation-resistant coating on Ti alloys and titanium aluminides (Subrahmanyam

and Annapurna, 1986; Umakoshi et al., 1989; Smialek and Humphrey, 1992), and as a wear-resistant surface layer (Uenishi et al., 1992). It should be noted that there is another Al-rich titanium aluminide, Al2Ti, which shows a higher resistance to cracking than identically processed Al3Ti with a similar hardness (Benci et al., 1993). Recently, Al3Ti has been applied as thinfilm interconnections in microelectronic devices where the Al3Ti acts as a diffusion barrier (Colgan, 1990; Gupta et al., 1993). Compared with Al metallization, intermetallic aluminides are advantageous because of their lower diffusion coefficients, and the best properties have been shown by Al3Ti films. Al^V

A13V is regarded as promising for applications in nuclear reactor technology and has been studied with respect to ductility improvements (Umakoshi etal., 1988). It is isostructural with Al3Ti and both form a continuous series of mixed crystals. The melting temperature is only slightly higher than that of Al3Ti and thus a similar deformation behavior is expected. The deformation mechanisms are indeed analogous to those in Al3Ti. However, the yield stress is higher than that of Al3Ti by a factor of about 2 and there is no ductility in compression (in vacuum) below about 400 °C. Microalloying with Al3Ti improves the ductility in compression - at the expense of the yield stress - to such an extent that A13(VO 95 Ti 0 05 ) shows a fracture strain of about 7%. Further alloying with Al3Ti leads to a deterioration in the properties, i.e. A13(VO 75 Ti 0 25 ) is significantly stronger than AI3V, but can only be deformed plastically at 700 °C and above. It is noted that according to a theoretical study with first-principles calculations, the

11.3 Titanium Aluminides and Related Phases

phase Al 3 Ru also has the D0 2 2 structure and is closely related to A13V (Paxton and Pettifor, 1992). The deformation behavior of Al 3 Ru is expected to be similar to that of A13V but with enhanced ductility. Al3Nb Al 3 Nb is a line compound and thus it usually contains second phases - in particular excess Al on the grain boundaries. It deforms by dislocation movements and by twinning, and it is brittle at room temperature (Shechtman and Jacobson, 1975). In spite of its room-temperature brittleness, Al 3 Nb is regarded as a candidate phase for high temperature structural service (Rodriguez et al., 1991). The deformation behavior, as well as the physical properties, of Al 3 Nb and Al3Nb-base alloys have been studied in more detail (Reip, 1991; Reip and Sauthoff, 1993); preliminary results have been presented in Sauthoff (1990 a, b), and the main findings are briefly reported in the following. In spite of the high density of Nb, the DO2 2 phase Al 3 Nb is still a lightweight material with a density of 4.54 g/cm3, and the mean thermal expansion coefficient is about 9.6 x l O ^ K " 1 between room temperature and 1000 °C which is lower than that for superalloys and steels. It melts congruently, in contrast to Al3Ti and A13V (Massalski et al., 1990), which makes preparation easier. Its melting temperature of 1605 °C is significantly higher than those of the other D0 2 2 phases, and thus a much higher strength - in particular at high temperatures - is expected. The Young's modulus, 246 GN/m 2 at room temperature and 224 GN/m 2 at 600 °C with a linear temperature dependence in between, is higher than those of the titanium aluminides and even higher than those of the superalloys (Table 11-2). Poisson's ratio has been

679

found to be 0.17, and similarly low values have also been found for Al3Ti and other intermetallics with complex crystal structures (Fleischer etal., 1989 b; Nakamura, 1991). The yield stress, i.e. the 0.2% proof stress of Al 3 Nb in air, is shown in Fig. 1119 as a function of temperature above 850 °C. Between 600 °C and 800 °C, Al 3 Nb fails catastrophically in air tests because of grain boundary oxidation, which is known as a pest phenomenon. Below 500 °C, the fracture strain in compression is smaller than 0.2 %, whereas plastic deformation in bending has only been observed above 1050 °C. The plane strain fracture toughness is only about 2 MN/m 3/2 at room temperature, as was also found by others (Schneibel et al., 1988). It is noted that the pest phenomenon is most pronounced for Al 3 Nb at a critical temperature of about 720 °C and under oxygen partial pressures between 10~ 15 and 10" 1 0 bar, whereas no pest is observed in pure oxygen (Steinhorst and Grabke, 1990; Grabke etal., 1991a). The pest-like disintegration of Al 3 Nb is proposed to be due to the selective oxidation of Al at the grain boundaries, which leads to Al depletion and the formation of Al 2 Nb (Tolpygo and Grabke, 1993). Other intermetallics suffer the pest phenomenon, too, and the mechanism is not well understood (Westbrook and Wood, 1964; Aitken, 1967; Meier and Pettit, 1992). The observed dislocation slip systems do not differ from those of the other D0 22 phases, and again twinning, which does not affect the order, is a major deformation mode at low and high temperatures. The number of independent deformation modes is smaller than prescribed by the Von Mises criterion, which contributes to the observed brittleness (see Sec. 11.2.3). The ductility at high temperatures results

680

11 Intermetallics

from thermally activated dislocation motions, i.e. creep. The creep strength of Al 3 Nb is comparatively low - a stress of 10 MN/m 2 produces 1 % strain in only 500 h and fracture in 2300 h - whereas the yield stress compares favorably with the superalloys. This illustrates the fact that the difference between the yield stress and the creep strength is much more pronounced for intermetallics than for conventional alloys. Creep of Al 3 Nb is controlled by dislocation climb which is accompanied by subgrain formation. The observed creep behavior corresponds to that of conventional disordered alloys and the creep rates are described by the known constitutive equations. This will be discussed in more detail with respect to NiAl (Sec. 11.4.3). The secondary creep rate follows the power law, i.e. Dorn equation for dislocation creep DGb

(11-2)

where 8 is the secondary strain rate, A a dimensionless factor, D the effective diffusion coefficient, G the shear modulus, b the Burgers vector, k Boltzmann's constant, T the temperature, a the applied stress, and the exponent n is between 3 and 5. For Al 3 Nb a stress exponent of 3.6 and an apparent activation energy of 360 kJ/mol was found, of which 35 kJ/mol are due to the temperature dependence of the shear modulus and the remaining 325 kJ/mol describe the temperature dependence of the other factors in the Dorn equation, in particular the diffusion coefficient. Diffusion data are scarce (Ogurtani, 1972; Slama and Vignes, 1972) and are not sufficient for analysis of the creep data. As already mentioned in Sec. 11.3.3.1, Al 3 Nb dissolves Ti, i.e. it can be alloyed with Al3Ti to form a continuous series of mixed crystals with a D0 2 2 structure which

are no longer line compounds. Figure 1119 shows the 0.2% proof stress of the ternary phase Al 3 (Nb 0 75 Ti 0 25 ) as a function of temperature in air, and it can be seen that in contrast to binary Al 3 Nb the fracture strain in compression is already greater than 0.2% at 200 °C, and furthermore there are no indications of the pest phenomenon. However, in bending the brittle-to-ductile transition temperature is only sightly lower than that for Al 3 Nb. Small additions of Hf and Li give a further lowering of the brittle-to-ductile transition temperature, as was already found for Al3Ti (see Sec. 11.3.3.2). It is noted that, as in the case of Al3Ti (see Sec. 11.3.4), attempts to alloy Al 3 Nb with a third element have been tried in order to produce the LI 2 structure instead of the D0 22 structure since the L l 2 structure is regarded as advantageous for developing ductile intermetallics (Li et al., 1992; Subramanian and Simmons, 1991). However, these efforts have not been successful. The respective ternary phase diagrams, including old literature from Germany and the former Soviet Union, have been reviewed by Kumar (1990). A more significant improvement in the toughness and the brittle-to-ductile transition temperature is obtained by using Al 3 Nb as a constituent in an intermetallic two-phase or multiphase alloy, i.e. by combining Al 3 Nb with another, less brittle phase. In particular Al 3 Nb forms a stable equilibrium with the B2 phase NiAl, i.e. intermetallic NiAl-Al 3 Nb alloys can be formed and are discussed in Sec. 11.4.3. Other Al3Nb-base alloys with a complex microstructure are produced by using novel processing methods including rapid solidification and powder metallurgy methods (Bowman and Noebe, 1989; Bowden, 1989; Lu et al., 1990; Alman and Stoloff, 1991; Rodriguez etal., 1991; Ayer and

11.3 Titanium Aiuminides and Related Phases

Ray, 1991; Ho and Sekhar, 1991). For example, a multiphase Al3Nb-base alloy which is dispersion-strengthened by TiB 2 was prepared by melt-spinning, pulverization and hot isostatic pressing (Ray and Ayer, 1992). Here much more work is necessary for estimating the potential of Al3Nb-base alloys for applications as structural, high-temperature materials. Apart from the mechanical behavior, a sufficient oxidation resistance is a prerequisite for any high-temperature applications. In spite of the fact that Al 3 Nb is being considered as an oxidation resistant coating for Nb-based alloys (Bowden, 1989), the oxidation resistance of bulk Al 3 Nb by scale formation is only limited (Perkins etal., 1988; Hebsur etal., 1989; Korinko and Duquette, 1990; Steinhorst, 1989; Steinhorst and Grabke, 1989; Grabke et al., 1990,1991 b). The oxidation resistance is improved by excess Al and by small additions of Hf (Steinhorst, 1989; Grabke etal., 1990, 1991b). Likewise Al 3 Nb that is macroalloyed with Cr and microalloyed with Y exhibits excellent oxidation resistance at 1200 °C, but it still suffers from severe environmental attack during deformation at intermediate temperatures (Raj et al., 1992a). Such alloys have also been studied with respect to strength and fracture toughness at low and high temperatures. Apart from these improvements by single-phase alloying, the oxidation resistance of Al 3 Nb is increased significantly by alloying with the already mentioned NiAl to form intermetallic multiphase alloys (Steinhorst, 1989; Grabke etal., 1990, 1991b), i.e. the multiphase Al 3 Nb-NiAl alloys are advantageous with respect to both their deformation behavior and their oxidation resistance, and they do not show the pest phenomenon.

681

Al3Ta Al3Ta with a D0 22 structure is closely related to Al 3 Nb (Bauer, 1939). The AlTa and Al-Ta-Ti phase diagrams have been studied recently (Subramanian et al., 1990b; Sridharan and Nowotny, 1983). Al3Ta has a still higher melting temperature (1627 °C) than Al 3 Nb and its deformation behavior seems to be similar (Pak etal., 1990). Fabrication of this phase is difficult because of the widely differing melting points of the constitutive elements. Nevertheless, it is being considered for composite materials developments (Alman and Stoloff, 1991; Anton, 1988; Shah etal., 1990; Kumar, 1991). Ni3V The above presented D0 2 2 trialuminides are line compounds and their ordered crystal structure is stable up to the melting point, which may be regarded as an indication of a strong ordering tendency. In contrast to these phases, the D0 2 2 phase Ni3V shows a range of homogeneity and its ordered crystal structure is only stable at lower temperatures, i.e. Ni 3 V disorders at 1045 °C to form an f.c.c. Ni-V solid solution (Massalski etal., 1990), which indicates a weak ordering tendency. The deformation mechanisms have been studied and analyzed in detail (Moreen etal., 1971; Vanderschaeve et al., 1979; Faress and Vanderschaeve, 1987; Yamaguchi and Umakoshi, 1990; Khantha et al., 1992). As in the case of the trialuminides, Ni3V deforms by twinning, which does not affect the order, and its yield stress is comparatively high, e.g. 720 MN/m 2 for polycrystalline Ni3V with large ordered domains and 1320 MN/m 2 for Ni3V with small domains. In contrast to the trialuminides, it shows considerable ductility in compression at room temperature. However, the

682

11 Intermetallics

ductility in tension is still small, i.e. Ni 3 V is also regarded as brittle (Liu and Inouye, 1979). Alloying Ni 3 V with Fe and Co leads to the phases (Fe,Ni)3V and (Fe,Co,Ni) 3 V with an L l 2 structure, which is the "parent" structure of the D0 2 2 structure. These ternary and quaternary phases are highly ductile with a tensile elongation of 40 % or more at room-temperature, and have been the subject of successful materials developments (Liu, 1984) which are described in Sec. 11.4.2 together with other L l 2 phases in connection with Ni 3 Al. This transition from the D0 2 2 structure to the L l 2 structure by alloying illustrates the fact that these closely related crystal structures (see Sec. 11.3.3.1) differ little with respect to energy and stability, as has been shown theoretically by ab initio calculations (Stocks e t a l , 1987; Carlsson, 1991; Freeman et al., 1991). In this context Ni 3 Nb with a D0 2 2 structure should be mentioned, which is of practical importance for some Ni-base superalloys. These alloys are strengthened by precipitated metastable Ni 3 Nb particles and show excellent weldability (Oblak etal., 1974). 11.3.4 Trialuminides with the L l 2 Structure 11.3.4.1 Basic Properties and Phase Diagrams

The cubic LI 2 structure is more symmetric than the tetragonal D0 2 2 structure (see Fig. 11-1) and has a sufficient number of slip systems according to the Von Mises criterion, and thus it should also be more deformable (George et al., 1991 b). In particular, after the successful "ductilization" of Ni 3 Al and Ni 3 V (see Sees. 11.4.1 and 11.4.2) the LI 2 structure is regarded as

most advantageous and promising for developing structural intermetallics. It has long been known that there are Al-rich ternary intermetallics with an ordered f.c.c. LI 2 structure (Raman and Schubert, 1965). In the preceding sections (Sees. 11.3.3.1 and 11.3.3.2) it was noted that the L l 2 and D0 22 crystal structures are closely related and the enery differences are small. A third, closely related structure is the D0 2 3 structure, which is another long-period ordered cubic structure based on the Ll 2 structure (Bauer, 1939). The binary trialuminides and the Al-rich ternary phases with these closely related crystal structures have recently been reviewed from theoretical and practical points of view with respect to stabilization of the LI 2 structure by alloying, and the prospects of materials developments on the basis of such L l 2 phases (Kumar, 1990, 1993; Freeman etal., 1991; George etal., 1991b). The main results are summarized in the following. The group IIIA element Sc, as well as some rare earths and actinides, forms a stable, binary trialuminide with an L l 2 structure, whereas the group IVA and VA elements form trialuminides with the D0 2 2 structure, i.e., Al3Ti and Al3Hf at high temperature, and Al3V, Al 3 Nb, and Al3Ta, or with the D0 2 3 structure, Al3Zr and Al3Hf, at low temperature. Ternary L l 2 phases are obtained by alloying Al3Ti with Cr, Mn, Fe, Co, Ni, Rh, Pt, Pd, Cu, Ag, Au, or Zn, whereas the L l 2 structure was not obtained with V, Nb, or Mo (Kumar, 1990; Freeman etal., 1991; George etal., 1991b; Nakayama and Mabuchi, 1993). Likewise, the Ll 2 structure has been produced by alloying Al3Zr with Sc, V, Cr, Mn, Fe, Ni, Cu, or Zn, and by alloying Al3Hf with Cu, whereas such alloying has not led to L l 2 phases in the cases of A13V, Al 3 Nb, and Al3Ta. The different alloying

11.3 Titanium Aluminides and Related Phases

behavior reflects the different stabilities of the competing structures and can be understood in view of the electron density distributions, as obtained by quantum mechanical ab initio calculations. This has been shown in particular for Al 3 Nb-Ni, which is attractive because of its advantageous combination of high melting temperature and low density (Inoue et al., 1991 b). It is noted that metastable binary trialuminides, i.e. powders of Al3Ti, Al3Zr, and Al3Hf, with Ll 2 structures could be obtained by mechanical alloying, i.e. low-temperature processing (Schwarz et al., 1992). The phase Pd 3 Mn, which orders below 530 °C to form the D0 2 3 structure in the absence of hydrogen and the LI 2 structure in the presence of hydrogen with sufficient partial pressure (Sowers et al., 1992), should be mentioned in this context. The compositions of the ternary Al-rich LI 2 phases can yary within a small range of homogeneity, and are/given approximately by various kinds pf formulae, e.g. Al 66 Ti 25 M 9 , Al 22 Ti 8 M 3 ox Al 5 Ti 2 M with M = Cr, Mn, Fe, etc. - jWith analogous formulae for the Zr variants - depending on the alloying element. It is not clear in what way the atoms are distributed in the Ll 2 lattice. Since these compositions are near those of the binary trialuminides in the respective ternary phase diagrams, these ternary L l 2 phases are commonly addressed as ternary L l 2 trialuminides with a partial substitution of Al by the third M element according to (Al 1 _ x M x ) 3 Ti, though there are arguments against this view (Durlu etal., 1991; Durlu and Inal, 1992 a). The respective phase diagrams have been discussed (see Kumar, 1990; Mazdiyasni etal., 1989; Mikkola etal., 1991; Nic et al., 1991). In view of the common composition range of the various trialuminides with an L l 2 structure, a stability

683

criterion for the existence of the L l 2 structure has been deduced which is given by a specific number of valence electrons per atom of the alloy if the effect of atomic interaction on the electronic distribution is considered according to the Engel- Brewer theory (Durlu and Inal, 1992 b). 11.3.4.2 Microstructure and Mechanical Behavior

Most studies are centered on Al 66 Ti 25 M 9 and its mechanical behavior is discussed in detail by George et al. (1991 b). Note that the processing of these alloys is crucial since the cast alloys usually contain residual pores and second phase particles. The elimination of these artifacts by post-solidification treatments, i.e., heat treatments, hot working, hot extrusion, or forging, is difficult and alloys of sufficient quality have only recently been obtained. Besides ingot metallurgy, powder metallurgy and rapid solidification have been used. Elastic moduli have been measured and the Young's modulus has been found to be 200 GN/m 2 for Al 67 Ti 25 Ni 8 and 192 GN/m 2 for Al 67 Ti 25 Fe 8 (George etal., 1991b), i.e., it is between those of Ti3Al and TiAl on the one hand and Al3Ti and Al 3 Nb on the other, and is of the order of those of the superalloys (Table 11-2 and Sees. 11.3.3.1 and 11.3.3.2). The compressive yield strength of Al 66 Ti 25 M 9 with M = Fe, Cr, Mn or mixtures of these transition metals is of the order of 300 MN/m 2 between room temperature and about 800 °C, i.e. there is a strength plateau or a slight positive temperature dependence (Yamaguchi and Umakoshi, 1990; Kumar etal., 1991b; Kumar and Brown, 1992b; Wu etal., 1993; Brown etal., 1993). Below room temperature there is a steep yield stress increase with decreasing temperature and

684

11 Intermetallics

above 800 °C the material softens in the usual way. Much higher yield stresses can be obtained by variations in the composition and the microstructure (George et al., 1991 b), and in particular by the introduction of dispersoids to produce particulate composites (Kumar et al., 1991c; Kumar, 1991; Otsuki and Stoloff, 1992). The mechanical behavior has been analyzed as a function of the degree of long-range order (Winnicka and Varin, 1993). The grain size dependence of the yield stress follows a Hall-Petch relationship (Pu etal., 1992). The tensile yield strength corresponds to the compressive yield strength in general, however, there are some complexities which are not yet understood (Kumar etal., 1991a, b; George etal., 1991b). The tensile or flexural elongation is only a fraction of a percent at room temperature and increases slowly with temperature with another ductility minimum at about 500°C (Z. L. Wu etal., 1990a,b, 1991; Kumar and Brown, 1992 a, c, 1993). The ductility of the single crystals tested to date is not higher than that of polycrystals and the various variations of composition and microstructure have led to only small ductility improvements (George et al., 1991 b; Brown etal., 1993). Slip systems, dislocation mobilities and reactions, fracture modes, bend ductility and creep strength have been studied in detail (Yamaguchi and Umakoshi, 1990; George etal., 1991b; Schneibel etal., 1992a; Morris et al., 1993 b, c; Sizek and Gray, 1993; Miura and Watanabe, 1993; Wu and Pope, 1993). It has to be concluded that the L l 2 trialuminides of the Al 66 Ti 25 M 9 type are brittle and do not fulfill the initial expectations with respect to ductility. Nevertheless, the L l 2 trialuminides are promising because of their low density and there is a potential for aerospace applications though this potential has not been

realized to date (Dimiduk et al. 1991). For such applications a sufficient oxidation resistance is a necessity, and indeed a cyclic oxidation resistance has been found which varies from poor for Al 67 Ti 25 Mn 8 to excellent for Al 67 Ti 25 Cr 8 (Parfitt et al., 1991). The present work is directed at the optimization of the composition and microstructure of Al 66 Ti 25 M 9 alloys by careful processing with respect to strength, ductility, toughness, and oxidation resistance. Besides these studies on Al3Ti-base trialuminides, similar work has been started on L l 2 trialuminide alloys, which are based on the D0 23 phase Al 3 Zr and which again contain Cu, Mn, or Cr (Virk and Varin, 1992; Schwarz etal., 1992; Varin et al., 1993). According to first results the mechanical behavior seems to be similar to that of the Al3Ti-base alloys, and there is a strong, positive temperature dependence for the yield stress.

11.4 Nickel Aluminides and Related Phases 11.4.1 M3A1 11.4.1.1 Basic Properties and Phase Diagram Ni 3 Al is the most studied and best known intermetallic because it has been used as a strengthening phase in the superalloys for a long time, and because its ductility problems can be overcome, i.e. it can be "ductilized" by microalloying with boron (Aoki and Izumi, 1979; Liu and Koch, 1983; Aoki, 1990). This means that it can be produced and tested without major experimental difficulties, and thus it was selected preferentially in the past for studies on the behavior of intermetallics (see

11.4 Nickel Aluminides and Related Phases

Sec. 11.1.2). The knowledge obtained on the physical and mechanical metallurgy of Ni 3 Al and its alloys has been reviewed in an authoritative way by Stoloff (1989). The nickel aluminide Ni 3 Al - known as the y' phase - crystallizes with the cubic LI 2 structure (Cu3Au-type) which results from the f.c.c. structure by ordering (see Fig. 11-1). Deviations from stoichiometry are accommodated primarily by antisite defects (Lin and Sun, 1993). The density of Ni3Al is 7.50 g/cm3 (see Liu et al., 1990) and thus is only slightly lower than that of the superalloys (see Table 11-2) which, however, is still of interest. The elastic constants have been studied experimentally and theoretically by various authors (e.g. Davies and Stoloff, 1965; Dickson e t a l , 1969; Kayser and Stassis, 1969; Foiles and Daw, 1987; Wallow etal., 1987; Yoo and Fu, 1991, 1993; Yasuda etal., 1991a, 1992). Young's modulus of cast polycrystalline Ni3Al at room temperature is about the same as that of pure Ni with a weaker temperature dependence (Stoloff, 1989),

685

i.e. it is slightly smaller than that of the superalloys (see Table 11-2). Data for thermal expansion, thermal conductivity, the Seebeck effect, and electrical resistivity have been presented (Stoloff, 1989). The thermal diffusivity of Ni 3 Al has been studied recently (Archambault and Hazotte, 1993). Figure 11-20 shows the binary Ni-Al phase diagram which considers some new results. Ni 3 Al melts incongruently at about 1383 °C to form liquid Ni-Al and solid, disordered, f.c.c. Ni-Al peritectically (Hilpert e t a l , 1987; Verhoeven etal., 1991), which is in contrast to the widely accepted phase diagram in Massalski et al. (1990). In addition, there is a eutectic equilibrium at 1380°C with the B2 phase NiAl and liquid Ni-Al. Ni 3 Al is stable, i.e. ordered, up to the melting point according to most investigators (Stoloff, 1989), which may, however, only be true near the stoichiometric composition since there are indications that the critical temperature of ordering may be slightly lower than the

composition in wt.% Ni 0 10 20 30

40

50

60

70

80

90

1800

Figure 11-20. Binary Ni-Al phase diagram which is based on that by Massalski et al. (1990) and considers the new results of Hilpert et al. (1987) and Verhoeven et al. (1991) with respect to the Ni3Al equilibria. 0 AI

10

20

30

40

50

60

70

composition in at.% Ni

80

90

100 Ni

686

11 Intermetallics

melting temperature for Ni 3 Al with a high Ni excess (Bremer etal., 1988; Yavari etal., 1991; Ramesh etal., 1992). Ni3Al can dissolve further elements, in particular other transition metals, and ternary Ni-Al based phase diagrams have been studied by various investigators (e.g. Guard and Westbrook, 1959; Kornilov, 1960; DasGupta et al., 1984; Chakravorty and West, 1985, 1986; Chakravorty etal., 1985; Nash and Liang, 1985; Vincent et al., 1988; Hong etal., 1989; Enomoto etal., 1991; Lee and Nash, 1991 a). The direction of the Ni 3 Al lobes in the isothermal sections of the ternary phase diagrams indicate what sites are occupied by the alloying additions. According to this and the results of specific site occupation studies, Ni sites are occupied by Co, Pd, Pt, Cu, or Sc, and Al sites are occupied by Ti, Zr, Hf, V, Nb, Ta, Mo, W, Zn, Ga, In, Si, Ge, Sn, or Sb, whereas Cr, Mn, or Fe occupy both sites with a slight preference for the Al sites depending on the composition (Ochiai etal., 1984; Miller and Horton, 1987; Shindo et al., 1988; Enomoto and Harada, 1989; Chiba etal., 1991; Hono etal., 1992). The case of Hf demonstrates the difficulties of site determination since there is also further strong experimental evidence of Hf occupying Ni sites instead of Al sites (Bohn et al., 1987 a, b). The chemical activity of Hf in Ni 3 Al has recently been determined, according to which, however, Hf substitutes for Al in Ni 3 Al (Albers et al., 1992). According to computer modeling, the site occupancy depends on the composition, i.e. different types of deviations from stoichiometry may change the site occupation by a third element (Hosoda et al., 1991). Deviations from stoichiometry, which are possible on both sides of the stoichiometric composition, lead to constitutional defects which are important for the me-

chanical behavior at both low and high temperatures since at low temperatures such defects may act as dislocation obstacles and at high temperatures they may promote diffusion. (This will be discussed in some more detail with respect to the deformation behavior of the B2 phases in Sec. 11.4.3.) Diffusion in Ni 3 Al has been studied by few investigators - in particular Chou and Chou (1985) and Hoshino etal. (1988) and has been reviewed and discussed with respect to mechanisms and defects (Bakker, 1984; Wever et al., 1989; Stoloff, 1989). The constitutional defects are antistructure atoms on both sides of stoichiometry, i.e. Al on Ni sites and Ni on Al sites, and the concentration of constitutional, i.e. athermal, vacancies is very small. The vacancy content of 6 x 10" 4 at the melting temperature and the vacancy formation enthalpy of 1.60 eV correspond to the respective values for Ni, i.e. the vacancy behavior of Ni3Al is similar to that of pure metals (Schaefer et al., 1992). The diffusion of Ni in Ni 3 Al is not very different from that in pure Ni and at high temperatures it is insensitive to deviations from stoichiometry. The diffusion of Al in Ni3Al is less well studied because a tracer is not readily available. Defects may interact with dissolved third elements which affects diffusion. In particular vacancies interact with B which is needed for "ductilization", and this leads to a complex dependence of the Ni diffusion coefficient on the Al and B content of Ni 3 Al (Hoshino etal., 1988). Data for the diffusion of the third elements, Co, Cr, or Ti, in Ni 3 Al are available (Minamino etal., 1992). 11.4.1.2 Microstructure and Mechanical Behavior

The mechanical behavior of Ni 3 Al and Ni3Al-based alloys, including creep, fa-

11.4 Nickel Aluminides and Related Phases

tigue and fracture, and in particular the effects of microstructure and alloying with further elements have been reviewed exhaustively (Stoloff, 1989). Only the most important features, i.e. the anomalous temperature dependence of the flow stress and ductilization by the boron effect, are discussed in the following, which also considers more recent studies. Other recent studies concern creep, including inverse creep after the first normal primary creep stage (Hazzledine and Schneibel, 1989; Hayashi e t a l , 1991a, b; Hemker e t a l , 1991; Miura etal., 1991; Schneibel and Hazzledine, 1992; Wolfenstine et al., 1992; Hemker and Nix, 1993; Zhang and Lin, 1993), superplasticity (Valiev etal., 1991; Ochiai etal., 1991; Yang etal., 1992; and the chapter by Mukherjee in Volume 6 of MST), fatigue (Matuszyk et al., 1990; Glatzel and Feller-Kniepmeier, 1991; Gordon and Unni, 1991; Gieseke and Sikka, 1992), and fracture (Kawabata and Takasugi, 1991; Takasugi, 1991a; Yoo and Fu, 1991). The recrystallization behavior has been studied in more detail only recently (Gottstein etal., 1989, 1991; Cahn, 1991;

687

Inoue and Inakazu, 1991; Ball and Gottstein, 1993; Zhou etal., 1993; Jena etal., 1993). Anomalous Temperature Dependence of the Flow Stress The plastic flow of Ni3Al has been studied in much detail by many authors. The results have been reviewed and discussed intensively and the basic mechanisms are now well understood (Stoloff and Davies, 1966; Paidar et al., 1981; Liu and Stiegler, 1984; Pope and Ezz, 1984; Stoloff, 1984; Liu and White, 1985; Izumi, 1989; Suzuki etal., 1989; Liu etal., 1990; Yamaguchi and Umakoshi, 1990; Liu, 1993 b; Saada and Veyssiere, 1993 b; and the chapter by Umakoshi in Volume 6 of MST). The outstanding feature is the positive, i.e. anomalous, temperature dependence of the flow stress which is illustrated in Fig. 11-21. Between room temperature and about 700 °C the flow stress increases with increasing temperature to reach a maximum and only at higher temperatures does normal softening occur.

o 6at% Nb D

A

1000•

10.Bat YoTi + 2ar%Cr iO.5ar %Ti 2ar% Cr

0A

/ s

NI3AI

/ /

^

V

/° a

/

sooFigure 11-21. Temperature dependence of the flow stress for cast polycrystalline Ni3Al and Ni3Al with various alloying additions (Thornton et al, 1970).

o200

400

600

temperature in °C

800

1000

688

11 Intermetallics

As already discussed (Sauthoff, 1986), the anomaly results from the anisotropy of the energy and mobility of the superlattice screw dislocations which determine the plastic deformation of Ni 3 Al. The screw dislocations can split on {111} planes and on {010} planes. The splitting on {010} is favored energetically because the energy of the antiphase boundaries (APBs) between the partials is lower on {010}. However, a superdislocation on {010} is sessile because the dislocation cores of the partials spread outside the plane of the APB. On the other hand, the superdislocation on {111} with higher energy is glissile because the core spreading is confined to the slip plane. This glissile state is metastable since the partials must first be coalesced into a single dislocation before cross slip back to {010} can occur. Hence dislocations are generated primarily on {111} planes on loading, and slip is confined to {111} at low temperatures. With rising temperature cross slip to {010} becomes possible by thermal activation, by which the dislocations are immobilized (Kear-Wilsdorf mechanism). The sessile dislocations act as obstacles and give rise to rapid strengthening with increased flow stress. Thus it is clear that the anomalous temperature dependence of the flow stress is observed only with sufficient strain, and indeed the minimum strain for this observable flow stress anomaly was found to be about 10~5 (Thornton et al., 1970). At temperatures above the flow stress maximum, the immobilized dislocations are mobilized again by enhanced thermal activation which leads to the familiar softening. The details of these mechanisms and reactions have been and still are the subject of specific studies, and the theoretical description becomes more and more elaborate (Yoo etal., 1988; Suzuki etal., 1989; Pope, 1991; Yoo and Fu, 1991; Webb et al.,

1993a; Schoeck, 1993; Dimiduk etal., 1993; Hirsch, 1993; Molenat et al., 1993; Saada and Veyssiere, 1993 a, b; Khantha etal., 1993a, b). From this discussion it is celar that the flow stress anomaly of Ni 3 Al is a function of the binding forces between the atoms. Indeed the anomaly is exhibited by various LI 2 phases to different extents, and the differences are closely correlated with the differences in stability between the L l 2 structure and other possible structures, in particular the D0 1 9 and D0 2 2 structures (Mishima et al., 1985; Suzuki et al., 1989). In Ni3Si and Ni 3 Ge the anomaly is even more pronounced than in Ni 3 Al, whereas it is weaker in Co3Ti, Zr 3 Al, Fe 3 Ga, and Cu3Au, and there is a continuous transition from a positive temperature dependence to the familiar negative one in the LI 2 ternary phase (Ni, Fe) 3 Ge when going from Ni 3 Ge to Fe 3 Ge. Likewise the flow stress maximum of Ni 3 Al can be increased and shifted with respect to temperature by ternary alloying additions, as is exemplified in Fig. 11-21 (see also Lall etal., 1979). It has to be noted that such flow stress anomalies are also shown by other phases with different crystal structures, e.g. FeCo-2V (B2), CuZn (B2), TiAl (Ll 0 ), Fe3Al (D03), Mg 3 Cd (D0 19 ), Ni 3 V (D0 22 ), and Fe 2 B (C16), and various mechanisms have been proposed for these phases (Stoloff and Davies, 1964; Westbrook, 1965; Paufler, 1976; Haasen, 1983; Stoloff, 1984; Kawabata etal., 1985). The positive temperature dependence of the flow stress of Ni 3 Al can be obscured by other concurrent strengthening effects. For example, the low-temperature flow stress can be increased by reducing the grain size to such an extent that it becomes higher than the original flow stress maximum at high temperature, i.e. a flow stress maximum is no longer discernible (Schul-

11.4 Nickel Aluminides and Related Phases

son, 1985). An analogous effect is produced when C is dissolved in Ni 3 Al to gradually increase the low-temperature flow stress (Jung and Sauthoff, 1989 b; Wunnike-Sanders, 1993; Wunnike-Sanders and Sauthoff, 1994). Ductilization It is well known that only polycrystalline Ni3Al is brittle and fails by intergranular fracture whereas Ni 3 Al single crystals are highly ductile, and thus the brittleness of polycrystals is attributed to grain boundary weakness (Liu and Stiegler, 1984). Correspondingly, directionally solidified, columnar-grained Ni 3 Al has been found to be ductile (Hirano and Mawari, 1993). The brittleness of polycrystalline Ni 3 Al can be reduced by macroalloying with Co, Cu, or Pt which reduce the ordering tendency of the Ni3Al (Chiba et al., 1992). Alternatively, the polycrystal brittleness of Ni3Al can be removed by microalloying Al-deficient Ni 3 Al with boron (Aoki and Izumi 1979; Aoki, 1990). The tensile elongation of polycrystalline Ni3Al with 24 at.% Al, which has been recrystallized for 30min at 1000 °C, increases with increasing B content from about 0 % for 0 % B to about 44% for 0.025 wt.% B (Liu et al., 1985). A further increase in the B content produces a maximum tensile elongation of about 54 % for 0.1 wt.% B, after which the ductility gradually decreases with increasing B content. This ductilization effect of boron depends sensitively on the Al content (Liu et al., 1985; Aoki, 1990). An Al content of 24 at.% corresponds to the Al solubility limit of Ni 3 Al and cannot be lowered, i.e. a further reduction only produces the disordered y-Ni-Al phase (see Fig. 11-20). With higher Al contents the ductilization effect becomes smaller and above the stoichiometric composition of 25 at.% Al,

689

polycrystalline Ni3Al with B is as brittle as without B. This stoichiometry effect can be changed significantly by alloying with a third element (Aoki et al., 1993). It should be noted that the ductility increase is accompanied by a significant increase in the ultimate tensile strength, which results from rapid strain hardening during plastic deformation (Liu et al., 1985). This dramatic ductilization effect of boron has been and still is the subject of elaborate experimental and theoretical studies which, however, have not yet led to an agreement on the physical reasons for the boron effect, i.e. the mechanistic understanding of this ductilization effect is still unclear. This has been the subject of a set of papers recently, and the present state of knowledge and the controversial issues have been overviewed (Liu, 1991b). The important factors are briefly summarized in the following. It has been found that B strongly segregates to grain boundaries in Al-deficient Ni 3 Al in equilibrium, as well as Al, which affects the character and strength of the bonding at the interface. The segregation of both elements depends on the bulk Al content. Boron does not segregate to free surfaces. It has been suggested that this B segregation enhances the cohesive strength at the grain boundary thus preventing cracking along the grain boundary. Theoretical first-principles calculations indeed confirmed the beneficial effect of B on the cohesive strength. A recent study of the microscopic deformation behavior has shown that during deformation grain boundaries with boron emit dislocations to relieve stress concentrations, whereas at grain boundaries without boron such stress concentrations are relieved by the nucleation and propagation of cracks along the grain boundaries (T. C. Lee et al., 1992).

690

11 Intermetallics

It has been proposed by others that this B segregation produces disorder at grain boundaries, thus facilitating slip transmission across them. This view is supported by a Hall-Petch type analysis of flow stress data as a function of the grain size. There are observations on disordered grain boundary layers which indicate the formation of the y-Ni-Al phase at the grain boundaries. This is not improbable since the Al content of Al-deficient Ni 3 Al is near the solubility limit, and in the case of the LI 2 phase Cu 3 Au it has been found (Tichelaar et al., 1992) that the order-disorder transition is preceded by the formation of disordered layers at the interfaces. However, such disordered layers have also been found in Ni 3 Al without B, and have not been found in many ductile Ni 3 Al alloys with B (see also Lin et al., 1993; Sun and Lin, 1993). A careful investigation of the order-disorder transformation in Ni 3 Al has shown that the order-disorder transition temperature is slightly lower than the solidus temperature for Ni 3 Al with less than 23 at.% Al (Cahn et al., 1987). This leads to a more complex ordering process with domain formation which may be beneficial for ductility, whereas in the case of stoichiometric Ni 3 Al with an order-disorder transition temperature above the liquidus temperature ordering occurs directly without domain formation. This B segregation affects the segregation of impurities such as O, S, and P which embrittle conventional alloys. However, such impurities seem to be of little importance for Ni 3 Al since Ni 3 Al, and other intermetallics of high purity with clean grain boundaries are still brittle. The experimental evidence for the various proposals is not clear and is even conflicting in some cases (Horton and Liu, 1990; Liu, 1991b). It has to be concluded

that the brittleness of intermetallics is a rather complex function of the crystal structure, composition and microstructure. Obviously this complexity, which seems to be more severe than for conventional alloys, excludes a simple, singlecause explanation. In any case, the ductilization results from changes in the character and strength of the bonding of the atoms involved which may be clarified by adequate quantum-mechanical calculations, and which may be characterized in an approximative way by differences in atomic parameters such as valency, electronegativity, and atomic size. Furthermore, effects of corroding environments have to be considered, as will be discussed in Sec. 11.4.1.3. Experimental results have recently been presented which indicate environmental embrittlement as a major cause of the low-temperature brittleness of polycrystalline Ni 3 Al, which is suppressed by microalloying with B (George et al., 1992, 1993 a, b; Zhu et al., 1993). However, this environment-induced embrittlement should be diminished by reducing the temperature, but this has not been observed, i.e. this aspect of the ductilization effect is also not yet clear (Lee and White, 1993). 11.4.1.3 Environmental Effects The oxidation behavior of Ni 3 Al has been studied repeatedly. The Al content of Ni3Al is sufficient for the formation of a stable, protective, A12O3 scale only at temperatures above about 1200 °C (Pettit, 1967; Meier, 1989). At lower temperatures Al diffusion in Ni 3 Al to the surface is too slow to avoid Al depletion at the surface, and thus the first formed A12O3 is overgrown by mixed Ni-Al scales. The initial stages of Ni 3 Al oxidation have been studied in detail (Bobeth et al., 1992). The oxi-

11.4 Nickel Aluminides and Related Phases

dation reaction is coupled with vacancy production which leads to the formation of voids below the scale and finally to scale spallation. Scale adherence is improved by the addition of Cr (Pan et al., 1991) and by the known oxygen-active elements - in particular Ti, Zr, Hf and the rare earths (Cathcart, 1985; Taniguchi and Shibata, 1989; Krasovec et al., 1992). The oxidation resistance of such Ni3Al-base alloys has been found to be superior to that of the conventional high-temperature alloys 800 H and 617 (Brill and Klower, 1991). A similarly advantageous corrosion behavior has also been shown in carburizing and chloridizing atmospheres, whereas the sulfidation behavior is less advantageous (Natesan, 1988; Brill and Klower, 1991). Apart from corrosion reactions at the surface, exposure to air or oxygen affects the mechanical behavior of Ni 3 Al and various effects of environmental embrittlement have been observed. At elevated temperatures of about 800 °C the ductility of ductilized Ni3Al shows a minimum which is more pronounced in air than in vacuum, and which is significantly alleviated by the addition of Cr (Liu and McKamey, 1990; Liu e t a l , 1990; Matuszyk etal., 1990; Khan et al., 1990). This embrittlement has been assumed to result from the effect of gaseous oxygen on the crack growth, whereas more recent results indicate that precipitated particles which are formed by internal oxidation may be more relevant to the embrittlement (DeVan and Hippsley, 1989; Chuang and Pan, 1992). The beneficial effect of Cr has been attributed to modifications to the mechanical properties of the scale (Taniguchi and Shibata, 1987; Hippsley etal., 1990), though a complete understanding of the ductility controlling mechanisms has not yet been reached (Liu and McKamey, 1990; Takeyama and Liu, 1992; Chuang and Pan, 1992).

691

Environmental embrittlement has also been observed at room temperature - in particular in moist air - depending on the Ni 3 Al composition and the dopant content, and it was concluded that this embrittlement is caused by hydrogen formation (Masahashi etal., 1988; Liu, 1991a; Wan et al., 1992a; George et al., 1993 b, c). Correspondingly, stress corrosion cracking has been observed in aqueous environments, which is also caused by hydrogen (Stoloff, 1989; Ricker etal., 1990). The embrittling effect of hydrogen has been studied directly and has been attributed to a reduction in the cohesive strength of the grain boundaries (Bond et al., 1989; Takasugi, 1991 b; Wan et al., 1992b). Nevertheless, the physical understanding of the hydrogen embrittlement effect is far from clear (Izumi, 1989; Takasugi, 1991a; Liu, 1991a; Li and Chaki, 1993 a). 11.4.1.4 Applications Detection of the ductilizing effect of boron on Ni 3 Al (see above) was the starting point for successful materials developments based on Ni 3 Al which have resulted in a series of closely related Ni 3 Al alloys for high-temperature applications (Liu et al., 1990). These alloys - known as nickel aluminides or advanced aluminides generally contain B at levels below 500 ppm, and Al below the stoichiometric composition, in order to obtain ductility at room temperature, and in addition Hf, Zr, Ta, and Mo at levels of up to 5 at.% to improve the strength at high temperatures and Cr up to 10 at.% to enhance the ductility at intermediate temperatures between 400 °C and 900°C (Liu et al., 1990; Sikka, 1990; Alexander and Sikka, 1992). The alloys may contain second phases and in particular the Cr-containing alloys generally contain 5 to 15 % of the disordered,

692

11 Intermetallics

Ni-rich phase, depending on the Al content (Liu et al., 1990; Khan et al., 1990; Liu and Kumar, 1993). The Ni3Al alloys compare favorably with many superalloys with respect to their short-term mechanical behavior, i.e. yield strength, which is higher than that of IN713C at elevated temperatures, and ductility, which is generally 25 % to 40 % at temperatures up to 700 °C and 15 % to 30 % at up to 1000 °C in air (Liu et al., 1990). The fatigue resistance of Ni 3 Al alloys is higher than that of Ni-base superalloys at temperatures below the ductility minimum, i.e. below 500 °C (Stoloff, 1989; Liu et al., 1990; Matuszyk etal., 1990; Gordon and Unni, 1991). The creep resistance of Ni 3 Al alloys is comparable to that of many superalloys, but it is not as high as that of some advanced, single-crystal, Ni-base superalloys which are used for jet engine turbine blades (Liu et al., 1990; Khan et al., 1990). In view of these properties it may be stated that the Ni 3 Al alloys are similar to the closely related superalloys and enlarge the superalloy spectrum. However, they cannot compete with the advanced aerospace superalloys which are used in aircraft engines, and thus Ni 3 Al alloys are unlikely to displace superalloys in such applications (Liu et al., 1990; Dimiduk et al., 1992). Nevertheless, Ni3Al alloys are promising for less demanding applications such as gas, water, and steam turbines, aircraft fasteners, automotive components, tooling and permanent molds where good combinations of strength and resistance against fatigue, wear including erosion and cavitation, and oxidation are needed (Liu et al., 1990). The technologies for fabrication and processing have been developed, i.e. Ni 3 Al alloys can be melted by air-induction melting with acceptable quality and by vacu-

um-induction melting, vacuum-arc remelting, and electroslag remelting for better qualities, ingots up to 1500 kg and components have been cast with low or no porosity depending on the boundary conditions, hot working is possible by conventional hot forging, isothermal forging and hot extrusion, the alloys can be cold worked and property data are available (Sikka, 1988, 1989,1990; Sanders et al., 1991; Alexander and Sikka, 1992; Haubold etal., 1992). Most of the alloys can be welded with good quality if it is done with care (Bittence, 1987; Chen and Chen, 1988; Stoloff, 1989; Liu et al., 1990; Santella, 1993; Li and Chaki, 1993 b), and cutting is accomplished by high-speed abrasive wheels (Sikka, 1990). Besides ingot metallurgy, powder metallurgy methods have been applied successfully for producing Ni 3 Al alloys (Stoloff, 1989; Sampath etal., 1991; Liang and Lavernia, 1991; Withers etal., 1991). Over recent years the developed technologies have been transferred to various industrial companies for the production of heating element wires, wear parts, diesel engine applications and aircraft fasteners, i.e. Ni3Al alloys are on their way to commercialization (Sikka, 1990; Liu and Kumar, 1993; Liu, 1993 a). An example is given by the ongoing development of an Ni 3 Al turbocharger rotor for heavy duty diesel engines, and it is expected that with further property improvements power cylinder components in high performance, state-of-the-art, diesel engines, i.e. valves, valve seats, piston crowns, piston rings, cylinder liners, and cylinder heads are likely candidates for future applications (Patten, 1990). The only application of Ni 3 Al alloys which has been realized up to now is their use in dies for hot pressing permanent magnetic alloys which are already on the market (Baker and George, 1992). Obvi-

11.4 Nickel Aluminides and Related Phases

ously, the transfer process from the laboratory to large scale production is handicapped by the limited deformability at elevated temperatures, which makes hot working difficult. Besides these developments, which are directed at specific applications, Ni3Al alloys are used for the development of intermetallic matrix composites which contain reinforcing particles or fibers of borides, carbides, oxides or carbon (Fuchs, 1989; Lee et al., 1990; Tortorelli et al., 1990; Alman and Stoloff, 1991; Kumar, 1991; McKamey and Carmichael, 1991; Mukherjee and Khanra, 1991; Brennan etal., 1992). Apart from the mechanical properties and the necessary corrosion resistance, the chemical compatibility of the used phases is of primary importance with respect to the long-term stability. It was found that SiC, B4C, and TiB2 react extensively with Ni3Al alloys, whereas very little reaction has been observed with A12O3 or TiC in Ni3Al (Fuchs, 1989; Lee etal., 1990; Brennan etal., 1992). It should be noted that Ni3Al alloys are used not only as the matrix material, but also as a reinforcing phase in, e.g. an Al alloy to form a metal-matrix composite (Metelnick and Varin, 1991). 11.4.2 Other Ll 2 Phases 11.4.2.1 General Remarks Apart from Ni 3 Al, Ni forms other intermetallic phases with an L l 2 structure and the composition Ni 3 X, where X is either an element near Al in the periodic table, e.g., Ni3Si, Ni 3 Ga, or Ni 3 Ge, or another transition metal near Ni, e.g. Ni 3 Fe or Ni 3 Mn; and there are also analogous examples for other transition metals, e.g. Fe 3 Ga, Fe 3 Ge, or Co3Ti (Villars and Calvert, 1991). LI 2 phases are among the first and best studied intermetallics because some of

693

them, e.g., Ni 3 Fe and Ni 3 Mn, as well as the classical example Cu 3 Au, transform into disordered solid solutions at temperatures below the melting temperature which allows the effect of atomic ordering on the physical properties and mechanical behavior to be studied (Sachs and Weerts, 1931; Dahl, 1936; Vidoz etal., 1963). Such phases with a critical temperature of ordering below the melting temperature are known as Kurnakov phases. Lately, the successful ductilization of Ni 3 Al and the subsequent materials developments have excited an enormous interest in LI 2 phases, which have been studied in great number and much detail in order to understand the physical mechanisms which control the plasticity (see, e.g., Veyssiere, 1992). In particular Cu 3 Au and Ni 3 Fe, which are experimentally easy, are used as model phases for detailed investigations of the microstructural features of deformation (e.g., Korner, 1991; Tichelaar and Schapink, 1991). Jerky flow has been observed in polycrystalline Cu 3 Au and Ni 3 Fe as well as in the L l 2 phases Zr 3 Al, Ni 3 Mn, and Ni3Al (Schulson, 1984). Ni 3 Ge has been used as a model phase for studying the effects of cyclic deformation in comparison to monotonic deformation in LI 2 phases (Pak etal., 1986; Inoue etal., 1991a). LI 2 phases have been studied systematically in comparison to Ni 3 Al in order to clarify the effects of changes in composition on the mechanical behavior. It has been known for a long time that polycrystalline Cu 3 Au and Ni 3 Fe are ductile at room temperature (Dahl, 1936; Vidoz et al., 1963) whereas polycrystalline Ni3Al is brittle, as was discussed in the preceding section. This difference in the ductility may be attributed to a difference in the ordering energy, since Cu 3 Au and Ni 3 Fe as Kurnakov phases disorder below the melting

694

11 Intermetallics

temperature and Ni3Al is ordered up to the melting point. However, there are other LI 2 phases that are ordered up to the melting point and which are not brittle, and it can be shown that the ductility of polycrystalline LI 2 phases is strongly correlated with the absolute valency difference of the constituent elements, which is reflected by the respective positions and distances in the periodic table; Ni3Si, Ni 3 Ga, Ni 3 Ge, Fe 3 Ga, and Ni 3 Al are brittle whereas Ni 3 Fe, Ni 3 Mn, Co 3 Ti, Cu 3 Pd and Cu 3 Au are ductile (Takasugi and Izumi, 1985d). Besides ductility, the temperature dependence of the flow stress may change significantly with changing composition. In particular, some L l 2 phases show an anomalous, positive, temperature dependence for the flow stress like Ni 3 Al, whereas other LI 2 phases show the conventional negative temperature dependence, as was already indicated in Sec. 11.4.1.2. In Ni3Si and Ni 3 Ge this anomaly is even more pronounced than in Ni 3 Al, whereas it is weaker in Co 3 Ti, Zr 3 Al, Fe 3 Ga, and Cu 3 Au, and there is a continuous transition from a positive temperature dependence to the familiar negative one in the L l 2 ternary phase (Ni,Fe) 3 Ge when going from Ni 3 Ge to Fe 3 Ge. This variation in the temperature dependence of the flow stress has been the subject of systematic studies and it can be shown that the differences are closely correlated with the differences in stability between the L l 2 structure and other possible structures, in particular the D0 1 9 and D0 22 structures (Mishima etal., 1985; Suzuki et al., 1989). Furthermore, there is a correlation between the anomalous temperature dependence of the flow stress and the elastic anisotropy which has been revealed by a systematic study of the elastic moduli of various L l 2 phases (Yasuda et al., 1991 a, b, 1992), and which supports

a theoretical model for the flow stress behavior (Yoo etal., 1988; Yoo and Fu, 1991). Apart from this scientific interest in LI 2 phases, the L l 2 structure is regarded as most promising for developing ductile intermetallic materials because of the successful "ductilization" of Ni 3 Al. Thus L l 2 phases have been the subject of many studies and indeed some L l 2 phases are of interest with respect to applications as structural materials. Examples are briefly discussed in the following section. Furthermore, other phases, which do not have an LI 2 structure, have been alloyed with further elements in order to produce the L l 2 structure. An example of this is given by the so-called trialuminides with an L l 2 structure, which result from alloying the binary Al3Ti phase with a tetragonal D0 22 structure with the transition metals Cr, Mn, Fe, Ni, and Cu, and which have been discussed in Sec. 11.3.4. However, it has to be noted that the expected ductilization of these trialuminides has not yet been achieved. Another example of obtaining the LI 2 structure by proper alloying is the multinary phase (Ni,Co,Fe)3V, on which a successful materials development has been based and which is also discussed in the following section. 11.4.2.2 Ll 2 Phases of Particular Interest M3Fe Ni 3 Fe is one of the ductile L l 2 phases, as mentioned in the preceding section. It shows an order-disorder transition at intermediate temperatures which has been studied in detail with respect to energetics and kinetics (Marty etal., 1990; Cahn, 1992; see also the chapter by Inden and Pitsch in Volume 5 of MST). The mechanical behavior is well known and the characteristics have been discussed with re-

11.4 Nickel Aluminides and Related Phases

spect to strength, dislocation behavior and fracture (see, e.g. Vidoz et al., 1963; Davies, 1963 b; Takasugi and Izumi, 1985 d, 1992; Izumi, 1989; Korner, 1991; D. G. Morris, 1992; Veyssiere, 1992). The elastic constants have been measured (Yasuda et al., 1992). As for other L l 2 phases, Ni 3 Fe is subject to hydrogen embrittlement at room temperature (Liu and Stoloff, 1993). The importance of Ni 3 Fe results from its advantageous physical properties, i.e. it is ferromagnetic with a high permeability (Kouvel, 1967; Dietrich, 1990). Thus Ni 3 Fe gave rise to the developments of the magnetically soft, high-nickel alloys which are known as Permalloys. Alloying with further elements improves the magnetic behavior as is demonstrated by the MolyPermalloys which typically contain 80% Ni and 4 to 5 % Mo - the balance being Fe - and the Mu-Metals which typically contain 77 % Ni, 5 % Cu and 2 % Cr, with the balance being Fe (Dietrich, 1990). Co3Ti A promising LI 2 phase is Co 3 Ti which is deformable over a wide temperature range (including room temperature) and has a high strength at elevated temperatures (Takasugi and Izumi, 1985 c; Takasugi etal., 1990b). However, there are indications of environmental embrittlement in air at ambient temperatures due to the embrittling effect of dissolved hydrogen (Takasugi and Izumi, 1986; Liu et al., 1989 b; Takasugi et al., 1990 b; Liu, 1991 a). The deformation behavior of Co3Ti has been studied in detail with respect to strength and ductility, and in particular an anomalous temperature dependence of the flow stress has been found, as in the case of Ni3Al (Takasugi and Izumi, 1985 b; Takasugi et al., 1987). In view of high temperature deformation, the diffusion coefficient

695

of Co in Co 3 Ti has been measured and the constitutional defects have been studied which result from deviations from stoichiometry (Takasugi and Izumi, 1985 a; Nakajima et al., 1988,1991). Co 3 Ti can be alloyed with third elements and the site occupation of the various elements has been investigated (Liu et al., 1986; Takasugi et al., 1990a). Such alloying additions can be used for varying the mechanical behavior, i.e. strength, ductility, flow stress anomaly, and sensitivity against hydrogen embrittlement (Liu et al., 1989 a; Takasugi etal., 1990b; Hasegawa etal., 1993a).

M3Si A promising L l 2 phase is Ni3Si, the deformation behavior of which, including the anomalous temperature dependence of the flow stress, is analogous to that of Ni 3 Al. It can be alloyed with Ti to form the ternary L l 2 phase Ni3(Si,Ti), which is being considered as a candidate phase for applications which require a high-temperature structural material or a corrosion resistant material since its strength is very high in comparison to other LI 2 phases, its ductility is high over a wide temperature range (particularly at lower temperatures when alloyed with B), it has an excellent corrosion resistance in hot sulfuric acid and likewise an excellent oxidation resistance in air at high temperatures, and it can be produced with good properties by various processing methods (Takasugi and Yoshida, 1991 b; Baker et al., 1993; Ulvensoen et al., 1993). Various studies have been directed at the alloying behavior and the effects on the mechanical properties, including fracture (Takasugi etal., 1991b; Takasugi and Yoshida, 1991a; Takasugi, 1991a; Takasugi and Izumi, 1992; Khantha et al., 1991; T. Zhang et al., 1991). The elastic moduli

696

11 Intermetallics

have been determined (Yasuda et al., 1992), and the dislocation behavior at low and high temperatures has been studied in detail (Yoshida and Takasugi, 1991a, 1992; Takasugi and Yoshida, 1992, 1993). The yield stress anomaly, which had already been reported in 1952 (Lowrie, 1952), has been clarified recently on the basis of ab initio calculations of the elastic constants and shear fault energies (Fu et al., 1993 b). Ni3(Si,Ti) alloys can be deformed superplastically (Nieh and Oliver, 1989; Takasugi et al., 1991a, c; Stoner etal., 1992) and environmental embrittlement has been observed (Liu and Oliver, 1991; Takasugi etal., 1991 d; Takasugi, 1991c). Recently, Ni3Si has been alloyed with Co 3 Ti to form two-phase intermetallic alloys with superior creep resistance (Hasegawa et al., 1993 b). Finally, it is noted that at very high temperatures an intermediate, brittle, monoclinic Ni3Si phase is stable, which may form the matrix of respective Ni-Si alloys with precipitated, ductile, Ni-rich solid solution and L l 2 Ni3Si as the toughening second phases (Baker et al., 1993; Li and Schulson, 1993). Zr3Al The LI 2 phase Zr 3 Al, which forms peritectically from Zr 2 Al and Zr, was studied extensively for use as the cladding material for water-cooled, nuclear power reactors since it has a low cross section for the absorption of thermal neutrons (Schulson, 1984; Liu et al., 1990; Parameswaran et al., 1990). In particular, the mechanical behavior was the subject of detailed studies (Schulson, 1984) according to which the strength is high with an anomalous temperature dependence for the flow stress. The work hardening rate exceeds that for any disordered or ordered f.c.c. alloy. Zr 3 Al, with polished surfaces, has a ductil-

ity of about 30% at room temperature with, however, a high notch sensitivity. Fast-neutron irradiation at low doses suppresses the notch sensitivity, whereas irradiation at higher doses produces an embrittling crystalline to amorphous phase transition which is accompanied by significant swelling. These irradiation effects precluded the use of Zr3Al in nuclear reactors (Parameswaran etal., 1990). Working of Zr3Al is limited to an area reduction of about 30 % and hot working is best done above the peritectic temperature, i.e. as two-phase Zr-Zr 2 Al, which is then followed by an annealing treatment below the peritectic temperature to produce the Zr 3 Al phase. The material has good oxidation resistance and can easily be fabricated into strip, rod, and tubing. (Fe,Co,Ni)3V The brittle phases Ni3V with a tetragonal DO2 2 structure and Co3V with a complex hexagonal structure can be alloyed with Fe in such a way that ternary (Ni,Fe)3V and (Co,Fe)3V and quaternary (Fe,Co,Ni)3V phases are formed which have the cubic LI 2 structure, and are ductile with tensile ductilities of about 40 % or more (Liu and Inouye, 1979; Liu, 1984; Liu et al., 1990). For these alloys the criterion for the stabilization of the L l 2 structure is a valence electron content of less than 7.89 per atom since higher valence electron contents between 7.89 and 8.6 are characteristic of Co3V and (Co,Ni)3V with hexagonal structures, whilst tetragonal Ni 3 V is characterized by 8.75 valence electrons per atom. The stability of these L l 2 alloys has also been studied by means of quantum-mechanical, first-principles calculations (Freeman et al., 1992). The LI 2 alloys (Ni,Fe)3V, (Co,Fe)3V, and (Fe,Co,Ni)3V disorder at a critical

11.4 Nickel Aluminides and Related Phases

temperature between 600 °C and 1000 °C, depending on the Fe content, to form disordered f.c.c. alloys. Recrystallization in the ordered state below the critical temperature is slow compared with the disordered state (Cahn, 1991), whereas diffusion does not seem to be affected much by the orderdisorder transition (Mantl etal., 1984). These L l 2 alloys can easily be fabricated because of their high ductility, they show an anomalous, positive temperature dependence for the flow stress like Ni 3 Al, strength decreases sharply above the critical temperature of disordering, and they have an excellent creep and fatigue resistance [for the latter see Ashok et al. (1983)] (Liu and Inouye, 1979; Liu, 1984; Liu et al., 1990). They can be welded (Bittence, 1987), but they are not sufficiently oxidation resistant above 600 °C because of the absence of Al (Liu, 1984), and they are subject to environmental embrittlement (Liu, 1991a; Nishimura and Liu, 1991, 1992a, b; Liu, 1992; Miura and Liu, 1992). In view of their advantageous mechanical behavior, they are regarded as promising for applications as structural materials in conventional power plants - in particular for steam turbines - as well as for nuclear power plants because irradiation produces relatively little swelling with, however, some lowering of ductility (Liu, 1984; Liu etal., 1990). 11.4.3 NiAl 11.4.3.1 Basic Properties NiAl is the best known example of the intermetallics with a cubic B2 structure (Fig. 11-1), which form one of the largest groups of intermetallics (Baker and Munroe, 1990). The physical and mechanical properties of NiAl have recently been reviewed in detail (Miracle, 1993). As is illustrated by the phase diagram in Fig. 11-20,

697

the phase NiAl has an extended range of homogeneity and melts congruently at about 1640 °C for the stoichiometric composition with 50 at.% Al. This melting point is higher than those of the constituent elements which indicates a strong bonding between Ni and Al and a corresponding high phase stability with a strong tendency for atomic ordering. This interpretation has been confirmed by quantummechanical, ab initio calculations according to which the strong Ni-Al bonding is of a mixed type, i.e. metallic with contributions of covalent and ionic bonding, which makes the Al atoms slightly electropositive (Fu and Yoo, 1992b; Lu et al., 1992). Indeed NiAl with a stoichiometric composition is highly ordered up to the melting point according to such ab initio calculations (Stocks et al., 1992). Deviations from stoichiometry result in constitutional disorder, i.e. constitutional point defects are produced which are antistructure atoms, i.e. excess Ni atoms on Al sites on the Nirich side of the stoichiometric composition and vacancies in the Ni sublattice on the Al-rich side, as has been shown experimentally and theoretically (Jacobi and Engell, 1971 a, b; Nakamura and Takamura, 1982; Koch and Koenig, 1986; Fu and Yoo, 1992a; Kogachi etal., 1992; S. M. Kim, 1992). In this way the number of valence electrons per alloy atom does not exceed a critical value for deviations from stoichiometry, in agreement with the HumeRothery rules for electronic compounds (Rusovic and Henig, 1980). These point defects interact with each other from which local ordering results (Georgopoulos and Cohen, 1981). Besides the constitutional point defects, there are the thermal point defects which again are vacancies in the Ni sublattice and Ni antistructure atoms in the Al sublattice since vacancies can only be created in the Ni sublattice,

698

11 Intermetallics

and any two such vacancies must be compensated by one Ni antistructure atom for a constant NiAl composition (Rusovic and Henig, 1980; Fu and Yoo? 1992a). The equilibrium configurations of the thermal vacancies have been studied theoretically (Kozubski, 1993). The density of NiAl is low at 5.9 g/cm3 for the stoichiometric composition compared with conventional Ni-base alloys, and it increases with decreasing Al content (Rusovic and Warlimont, 1977; Harmouche and Wolfenden, 1987). It is noted that the density change per unit of Al content is larger for Al-rich NiAl than for Nirich NiAl because of the difference in the defect structure. The elastic behavior of NiAl has been studied repeatedly and the elastic moduli have been determined experimentally for polycrystals and single crystals as a function of composition and temperature (Wasilewski, 1966; Rusovic and Warlimont, 1977, 1979; Harmouche and Wolfenden, 1985, 1987). The elastic behavior has also been studied theoretically by quantum-mechanical, ab initio calculations, and the resulting elastic moduli are in close agreement with the experimental values (Yoo et al. 1990; Fu and Yoo, 1992 b; Freeman et al., 1992; Yoo and Fu, 1993). The Young's modulus of polycrystalline NiAl with a stoichiometric composition is about 235 GPa at room temperature (Harmouche and Wolfenden, 1987). The elastic moduli are functions of the composition, and the Young's modulus reaches a maximum of slightly more than 235 GPa, not at the stoichiometric composition as is expected, but at about 48 at.% Al which may be related to the difference in the defect character on both sides of stoichiometry. The effect of excess vacancies, which are produced by quenching from high temperatures, on the elastic behavior was studied

by Rusovic and Henig (1980). The elastic anisotropy is larger than that of corresponding disordered metals with an A2 structure, e.g. b.c.c. Fe, and corresponds to that of f.c.c. Cu or Ni (Al structure) (Wasilewski, 1966). As a consequence of this anisotropy, the elastic constants of polycrystalline NiAl depend sensitively on texture. Ni-rich NiAl shows anomalies in its elastic behavior, and in particular there is an anomalous temperature dependence which indicates lattice softening and is related to the martensitic transformation (Rusovic and Warlimont, 1975,1977). The martensitic transformation of NiAl is discussed in the following section. Diffusion in NiAl has been studied repeatedly, from which a reliable data basis has resulted (see, e.g. Hagel, 1967; Hancock and McDonnell, 1971; Shankar and Seigle, 1978; Hao etal., 1985; Wever, 1992). The diffusion coefficient is expected to reach a minimum at the stoichiometric composition since any deviations from stoichiometry reduce the degree of atomic order by introducing constitutional point defects which enhance diffusion. Indeed the diffusion coefficient for Ni in NiAl reaches a minimum at about 49 at.% Al (Hancock and McDonnell, 1971; Shankar and Seigle, 1978), which is slightly below the stoichiometric composition and is related to the effects of point defects, i.e. vacancies and antisite atoms, on diffusion (Wang and Akbar, 1993). The tracer diffusion coefficient of Al in NiAl cannot be measured directly because the necessary Al isotope is not available. The atomic migration processes in NiAl are more complex than in disordered b.c.c. alloys because of the ordered atom distribution and, in spite of detailed analyses of the atom jump processes, a complete understanding and modeling of diffusion in NiAl has not yet been reached (Bakker, 1984; Wever et al., 1989;

11.4 Nickel Aluminides and Related Phases

Koiwa, 1992; Bakker et al., 1992). Only recently a model, which is based on a combination of two mechanisms, has been proposed for describing the composition dependence of diffusion in B2 phases (Kao and Chang, 1993). It should be noted that the understanding of the basic diffusion processes for other intermetallics is still less than for NiAl and other B2 phases (Wever, 1992). Diffusion studies of multicomponent systems are rare, and with respect to NiAl base alloys only data for the ternary phase (Ni,Fe)Al are available from a systematic study of the system Ni-Al-Fe (Cheng and Dayananda, 1979; Dayananda, 1992). Recently, the effect of Cr on diffusion in NiAl has been studied (Hopfeetal., 1993). 11.4.3.2 Phase Diagram and Martensitic Transformation

The equilibria of NiAl with neighboring nickel aluminides at high and low temperatures are shown by the Ni-Al phase diagram in Fig. 11-20. Al-rich NiAl forms an equilibrium with Al 3 Ni 2 only at higher temperatures, whereas at lower temperatures an equilibrium is formed with the phase Al 4 Ni 3 which has only recently been identified (Mukherjee etal., 1979; Ellner et ah, 1989). Its crystal structure is related to the B2 structure of NiAl and contains ordered vacancies. Ni-rich NiAl is in equilibrium with only Ni 3 Al at high temperatures and the mutual solubilities are shown in the phase diagram (according to Nishimura and Liu, 1992b; Bremer et al., 1988; Fang and Schulson, 1992). At low temperatures, Ni-rich NiAl forms an equilibrium with Ni 5 Al 3 which has been studied in detail (Robertson and Wayman, 1984; Khadikar and Vedula, 1987; Khadikar et al., 1993). In addition, precipitate particles of a metastable Ni 2 Al phase with either a trigonal or a monoclinic crystal

699

structure have been observed in Ni-rich NiAl after quenching and subsequent annealing at intermediate temperatures (Murthy and Goo, 1993; Muto etal., 1993 a, b). Ni-rich supersaturated NiAl, i.e. NiAl with more than 60 at.% Ni, can transform martensitically if decomposition and precipitation of Ni 3 Al and/or Ni 5 Al 3 can be avoided by annealing at high temperatures with rapid quenching to low temperatures. The martensite formation temperature Ms increases linearly with increasing Ni supersaturation from about -240°C for 60 at.% Ni to about 1000°C for 70 at.% Ni (Au and Wayman, 1972; Smialek and Hehemann, 1973; Ochiai and Ueno, 1988). The martensitic transformation produces the ordered, face-centered tetragonal L l 0 structure with a tetragonality of c/a = 0.86, and there is an internally twinned substructure (Chakravorty and Wayman, 1976 a, b). In some cases a hexagonal martensite with 2H stacking (a two-layer structure) was also observed (Litvinov et al., 1974; Martynov et al., 1983; Tanner et al., 1990), and furthermore a monoclinic structure with rhombohedral 7R stacking (a seven-layer structure) was found depending on the Al content (Martynov et al., 1983; Tanner et al., 1990; Khachaturyan etal., 1992; Shimizu and Tadaki, 1992). The 7R structure is intermediate in the course of the martensitic transformation of NiAl and is derived from L l 0 by introducing periodic stacking faults. It can be shown theoretically that the energy barriers for the martensitic transition from the cubic B2 structure to the tetragonal L l 0 structure are lowered by the formation of the intermediate 7R phase (Yegorushkin et al., 1985; Stocks et al., 1992; Khachaturyan et al., 1992; Sluiter et al., 1992). The elastic moduli of the formed martensite have been determined (Robertson, 1990).

700

11 Intermetallics

The martensitic transformation of Nirich NiAl can be induced by external applied stresses; it is thermoelastic, superelasticity has been observed, and the martensite deforms by twinning and detwinning (Au and Wayman, 1972; Chakravorty and Wayman, 1976b; Enami etal., 1981; Murakami etal., 1992; Kim and Wayman, 1992). Thus all prerequisites for the shape memory effect are fulfilled and indeed the shape memory effect has been produced with NiAl irrespective of the method of alloy preparation - ingot metallurgy or powder metallurgy, i.e. it can be shown that the transformation characteristics do not depend on the processing route (Ochiai etal., 1987; Kim and Wayman, 1990). According to a computer simulation, the stress-induced martensitic transformation is also expected for NiAl which is not supersaturated with Ni, but which is subject to high stress concentrations, e.g. at crack tips (D. Kim et al., 1991). NiAl can be alloyed with further elements in order to form ternary phases with a B2 structure which is then known as an L2 0 structure, too, or to obtain additional phases in equilibrium with NiAl. Fe, as well as Co, can substitute for Ni in NiAl completely without affecting the B2 structure, as is expected in view of the binary B2 phases FeAl and CoAl. Correspondingly, the ternary Ni-Fe-Al phase diagram, which is of importance with respect to conventional high temperature alloys, shows the extended B2 phase field and the respective equilibria with Ni 3 Al and Al-rich phases on the one hand and with disordered b.c.c. Fe and f.c.c. Fe on the other (Bradley and Taylor, 1938; Dannohl, 1942; Bradley, 1951; Hao et al., 1984). The N i Al-Co system behaves in an analogous way (Hao et al., 1984; Ishida et al., 1991 a, 1993). Other ternary systems have been re-

viewed (Raman and Schubert, 1965; Kumar, 1990). Another important alloy system is the system Ni-Al-Ti which contains various ternary intermetallics besides NiAl, in particular the Heusler-type phase Ni2AlTi with an L2X structure which is derived from the B2 structure (see Sec. 11.4.5), and the hexagonal Laves phase Ti46Ni27Al27 with a C14 structure (see Sect. 11.8) (Raman and Schubert, 1965; Nash and Liang, 1985; Lee and Nash, 1991b; Yang etal., 1992b). Similar situations are found in the phase diagrams for Ni-Al-Nb (Benjamin et al., 1966; Reip, 1991; Ochiai etal., 1991) and Ni-Al-Ta (Paketal., 1988). Ternary alloying is of particular importance with respect to the martensitic transformation of NiAl since it offers the possibility of controlling the martensite formation temperature Ms besides other transformation features. Co reduces the transformation temperature, reduces the transformation strain anisotropy and increases the interfacial mobility (Russell et al., 1989). It is noted that the curves of constant Ms in the ternary Ni-Al-Co diagram remain in the two-phase field, i.e. only the supersaturated (Ni,Co)Al can transform martensitically. The effect of other elements, i.e. Ti, V, Nb, Ta, Cr, Mo, W, Fe, Mn, Cu, Si (Kainuma et al., 1992 a) and B (Xie etal., 1993), on Ms has been studied recently and it has been found that alloying with Fe leads to a situation which is similar to the Ni-Al-Co case, whereas alloying with Mn enlarges the B2 field in the ternary Ni-Al-Mn phase diagram to such an extent that the curves of constant M s are shifted into the B2 field, i.e. it is expected that supersaturation is not necessary for the martensitic transformation of (Ni,Mn)Al (Kainuma etal., 1992b,c).

11.4 Nickel Aluminides and Related Phases

701

11.4.3.3 Microstructure and Mechanical Behavior

Figure 11-22 illustrates the variation of the flow stress of polycrystalline NiAl with temperature and composition. The curves show the normal temperature dependence with high strength at low temperatures, a comparatively steep strength decrease at intermediate temperatures around half the melting temperature (about 680 °C for NiAl) and low strength with further softening at high temperatures. It can be seen that variations in the composition lead to variations in the strength, and these effects may be different at low and high temperatures with a transition range around half the melting temperature (Sauthoff, 1990b). The softening at high temperatures is related to thermally activated creep processes which are the subject of the next section. Deviations from stoichiometry produce constitutional defects which enhance diffusion and lead to softening at high temperatures, whereas at low temperatures these defects are immobile and act as strengthening deformation obstacles. These different effects of deviations from stoichiometry at low and high temperatures were studied in detail for binary NiAl (Vandervoort etal., 1966). Alloying of NiAl with a third element, e.g. Fe, leads to solid-solution strengthening in the total temperature range of Fig. 11-22. These effects depend on both the temperature and the strain rate, as was shown by Rudy and Sauthoff (1985) and Rudy (1986). The effects of various other solutes on the high-temperature strength have been studied (Vedula et al., 1985; Nathal, 1992; Cotton etal., 1993b). Besides composition, processing was found to be a strength controlling factor which underlines the important role of the mi-

0

300

600

900

1200

temperature in °C

Figure 11-22. Flow stress (0.2% proof stress in compression at 10~ 4 s~ 1 strain rate) as a function of temperature for binary and ternary NiAl phases, i.e., stoichiometric NiAl (circles), stoichiometric (Ni0 8 Fe 02 )Al (squares), and off-stoichiometric (Nix 0 Fe 0 2)Ai0 8 (triangles) (Rudy and Sauthoff, 1985; Rudy, 1986).

crostructure in strengthening. The highest strengths were obtained with alloys with an excess of solute which was precipitated as second phase particles. NiAl alloys with second phases are discussed in Sees. 11.4.3.4 and 11.4.3.6. Plastic deformation of NiAl occurs predominantly by the movement of perfect dislocations in a <100> slip direction which is the shortest lattice translation vector in the B2 structure, and the preferred slip planes are {Oil}, whereas slip on {001} is also possible (Ball and Smallman, 1966; Strutt and Dodd, 1970; Baker and Munroe, 1990; Darolia et al., 1992b; Y Zhang etal., 1993). The details of slip in single crystals have been studied recently as a function of crystal orientation, deformation mode and temperature (Takasugi et al., 1993 b, c). In view of the high ordering energy of NiAl (see Sec. 11.4.3.1), creation and movement of dislocations is expected to be difficult, i.e. high strength and low ductility are expected. In a more fundamental sense, it is expected that the high strength and low ductility are caused by

702

11 In term eta I lies

the strength and character of the atomic bonding. However, for NiAl there seems to be no simple correlation between the directionality and ionicity of the bonding and the mechanical behavior, and much more work is necessary (Schultz and Davenport, 1992). As a consequence of the available slip systems, the strength and ductility are highly anisotropic with a "hard" <100> direction and "soft" <110> and <111> directions (Darolia e t a l , 1992b; Glatzel etal. 1993b; Takasugi et al., 1993a). NiAl with the "hard" orientation shows practically no ductility - in spite of indications of local plastic deformation (Vehoff, 1992) - below the brittle-to-ductile transition temperature (BDTT) which is of the order of 350 °C - corresponding to 0.33 Tm (Tm = melting temperature) depending on the deformation rate - whereas NiAl with a soft orientation shows ductilities up to 2.5% below the BDTT, which is of the order of 200 °C, corresponding to 0.25 Tm (Lahrman etal., 1991). The low ductility of NiAl at low temperatures is proposed to result from low dislocation mobilities because of extended dislocation cores (Farkas etal., 1991; Mills etal., 1993; Kitano and Pollock, 1993). A recent study has shown the importance of a sufficient density of mobile dislocations since an increased fracture toughness has been obtained by avoiding any annealing treatments (Hack et al., 1992). At the BDTT the ductility increases rather steeply with increasing temperature due to thermal activation which, however, is not yet understood in detail (Darolia et a l , 1992b). Like strength and ductility, the fracture toughness is a sensitive function of the orientation and the temperature, and this has recently been studied in detail (Reuss and Vehoff, 1990, 1992; Vehoff, 1992; K.-M. Chang etal., 1991, 1992; Darolia etal.,

1993). Corresponding to the low ductility, the fracture toughness of NiAl at lower temperatures is low. However, it is even lower in the soft <110> direction than in the hard <100> direction which is assumed to be a result of preferred cleavage on {110} planes. However, the details of the fracture processes are more complicated and not yet fully understood, according to recent observations (Schneibel et al., 1993 a). The fracture toughness increases substantially above 200 °C, independent of the crystal orientation which is in contrast to the BDTT for the ductility. The crack tip stress and deformation fields in NiAl single crystals have been analyzed in detail (Saeedvafa and Rice, 1992). Polycrystalline NiAl is usually found to be brittle with practically no ductility (Baker and Munroe, 1990; Reuss and Vehoff, 1992). Crack propagation occurs in an instationary way, and the fracture resistance is not sensitive to the loading rate (Schneibel etal., 1993b). Only with carefully processed, high-purity NiAl with an exact stoichiometric composition can a tensile ductility of 2.5% be obtained at room temperature, whereas similarly processed off-stoichiometric NiAl is brittle (Hahn and Vedula, 1989; Baker etal., 1991). Obviously the constitutional defects in off-stoichiometric NiAl reduce the ductility. In any case, the brittle-to-ductile transition temperature can be reduced significantly by reducing the grain size (Schulson, 1985; Nagpal and Baker, 1990 b; Chan, 1990; Cheng, 1992). The limit to the grain size reduction is nanocrystalline NiAl which has been prepared successfully with a grain size of about 10 nm (Haubold et al., 1992). The recrystallization behavior of NiAl was studied by Haff and Schulson (1982). Apart from the grain size effect, the brittleness can be reduced by introducing additional dislocation sources. This effect

11.4 Nickel Aluminides and Related Phases

has been observed in the case of NiAl with oxide layers on the surface or with particles in the alloy where dislocations are emitted from the respective phase boundaries (Prakash and Pool, 1981; J. T. Kim et al., 1991; Gibala et al., 1993). As in other cases, a higher ductility and toughness is observed with a superimposed hydrostatic pressure (Margevicius etal., 1993 a, b; Pawelski etal., 1994). One of the reasons for the brittleness of polycrystalline NiAl is the insufficient number of slip systems since there are only three independent slip systems with the <100> slip vector, whereas at least five independent deformation systems are necessary for general homogeneous deformation according to the Von Mises criterion (von Mises, 1928; Baker and Munroe, 1990; Darolia etal., 1992 b). Besides <100>, other slip vectors, i.e. <110> and <111>, have been observed under special circumstances (Field etal., 1991a; Kim and Gibala, 1991; Miracle, 1991; Veyssiere and Noebe, 1992; Dollar et al., 1992). The <111 > dislocation which corresponds to the common slip direction in b.c.c. alloys is a superlattice dislocation in the B2 structure consisting of two superpartials with an antiphase boundary (APB) in between. The various dislocation types - in particular the dislocation cores - have recently been studied in detail by direct observations and computer simulations, and the consequences of the dislocation mobilities have been discussed (Mills and Miracle, 1993; Parthasarathy et al., 1993; Mei and Cooper, 1993). With these additional slip systems the Von Mises criterion can be fulfilled and an increased ductility is expected. It has been found experimentally that alloying NiAl with Cr, which reduces the APB energy, may indeed increase the propensity to < 111 > slip (Field et al.,

703

1991b) and produces a strong solid-solution strengthening effect (Cotton et al., 1993 a), though the conditions and mechanisms are still in discussion (Cotton et al., 1993 a, c). An improved ductility has been found for NiAl-Cr only in compression (Kowalski and Frommeyer, 1992), whereas the tensile ductility was actually reduced and the brittle-to-ductile transition temperature was increased (Field et al., 1991 b; Cotton etal., 1993a). Obviously the Von Mises criterion is not a sufficient criterion and the material fractures in a brittle manner because the fracture strength is lower than the yield strength. Other elements which have been reported to promote <111> slip are Mn, V, Ti, and Zr (George etal., 1991a; Field etal., 1991b). With Fe substituting Ni there is a continuous transition from <100> slip to <111> slip (Patrick etal., 1991). Quantum-mechanical, ab initio calculations have shown that the APB energy is indeed lowered by the addition of Cr, Mn, or V to NiAl (Hong and Freeman, 1989). Another reason for the brittleness of polycrystalline NiAl may be grain boundary weakness. It was found that alloying with a small amount of B, which is beneficial in the case of Ni 3 Al and segregates to the grain boundaries in NiAl, changes the fracture mode from intergranular fracture to transgranular fracture (George etal., 1991a; Xie etal., 1993). However, ductility and fracture toughness are not improved, and a higher amount of B even reduces the ductility which may be related to an increase in the yield strength to levels above the fracture strength and to the precipitation of borides (Wu and Sass, 1993 a, b; Jayaram and Miller, 1993; Schneibel et al., 1993 b). The hardening effect of B additions is observed for off-stoichiometric, Ni-rich NiAl and is due to the variation in the B solubility (Tan et al.,

704

11 Intermetallics

1993). C and Be do not segregate at grain boundaries and do not change the fracture mode. Nevertheless, Be improves the ductility of NiAl slightly and the reasons for this behavior are unclear. It is known that alloying with, e.g. Fe, can improve the ductility of NiAl (Baker and Munroe, 1990). Indeed substantial increases in the tensile ductility at room temperature of monocrystalline NiAl with a "soft" orientation have been obtained by alloying with small amounts, i.e. about 0.1 at.% of Mo, Ga. or Fe (Darolia et al., 1992a): 0.1 at.% Fe produces 6 % ductility in contrast to 1 % ductility without Fe whereas 0.5 % and more Fe produces only 2 % ductility. The physical reasons for this ductilization effect are not yet clear (Noebe and Behbehani, 1992). 11.4.3.4 Creep High temperature deformation of metallic and other materials is controlled by various deformation mechanisms depending on the temperature and the deformation rate, as is illustrated by the so-called deformation maps for the respective materials (Frost and Ashby, 1982). The comparison of such maps for various materials with different crystal structures and types of atomic bonding shows that the deformation behavior at temperatures above a third of the melting temperature is controlled by creep processes, not only at low deformation rates of10~ 7 s" 1 and lower, which are used during typical creep experiments, but also at higher rates of say 10" 3 s" 1 which are used during short-term tests for flow stress determinations, and this may also be assumed to be the case for intermetallic phases (Sauthoff, 1990 b). Thus any discussion of high temperature deformation behavior has to be centered on the strength-controlling creep mechanisms.

The creep behavior of NiAl has been studied and discussed repeatedly in the past (Stoloff and Davies, 1966; Vandervoort et al., 1966; Strutt and Dodd, 1970; Strutt and Kear, 1985; Yaney and Nix, 1988). In addition, the creep behavior of NiAl and NiAl-based alloys has been investigated systematically and the findings have been the subject of various summarizing reports (Jung etal., 1987; Sauthoff, 1990 b, 1991 a, b, 1992). Further creep studies have been directed at NiAl single crystals (Forbes et al., 1993; Glatzel et al., 1993 b; Forbes and Nix, 1993). In the following sections the main characteristics of the creep behavior of NiAl are reviewed, the effects of structure and composition changes are outlined and both single-phase materials and alloys which contain other intermetallic phases besides the B2 phases are considered. The creep behavior of NiAl is discussed here in further detail since the obtained understanding of the rate-controlling mechanisms is also valid for other intermetallics, and thus the NiAl case may serve as a characteristic example for other intermetallic alloy systems. Creep Resistance of Single-Phase B2 Alloys The creep behavior of the ternary B2 phase (Ni,Fe)Al was studied in detail as a function of stress, temperature, composition, and grain size (Rudy and Sauthoff, 1985; Rudy, 1986; Jung et al., 1987). At high temperatures, e.g. 60 % of the melting temperature or higher, the secondary creep at rates between about 10 s and 10V - 6 s" 1 exhibits power law behavior, i.e. the observed secondary creep rates are described by the familiar Dorn equation for dislocation creep [Eq. (11-2)] (Mukherjee etal., 1969). Dislocation creep of conventional disordered alloys is produced by gliding and

11.4 Nickel Aluminides and Related Phases

climbing dislocations. If climb is the slower step, as in pure metals, the creep rate is controlled by dislocation climb, which gives rise to a well defined subgrain structure, and the stress exponent is 4 or 5. This is characteristic of class II alloys. Otherwise, viscous dislocation glide is rate controlling, which leads to dislocation tangles without subgrain formation and a stress exponent of 3, with behavior characteristic of class I alloys (Sherby and Burke, 1967; Nix and Ilschner, 1979). The secondary creep behavior of (Ni,Fe)Al shows analogous characteristics. In the Ni rich phases and in the binary NiAl a well-defined substructure is found after creep. The subgrain size is of the order of 10 jam, and the dislocation density within the subgrains is about 10 1 1 m~ 2 . In agreement with this, stress exponents between 4 and 4.5 have been found for the Ni-rich phases, i.e., these phases behave like class II alloys with dislocation climb controlling the creep. In the Fe-rich phases and in FeAl, however, no subgrain formation has been observed even after long creep times. The dislocation density remains high at about 1014 m~ 2 , and the stress exponent varies between 3 and 3.6. This indicates class I behavior, i.e. here the creep is controlled by the viscous glide of dislocations. In both cases only undissociated <100> dislocations have been observed. Obviously the driving force and the atom mobility which are necessary for subgrain formation are sufficient only in the Ni-rich phases. The line energies, line tensions and derivatives of dislocations have been calculated theoretically for NiAl, and good agreement with observed dislocation configurations after creep has been found (Glatzel et al., 1993 a). These findings show that the dislocation creep of such intermetallic alloys and

705

of conventional disordered Ni-base or Fe-base alloys is controlled by the same mechanisms. Even the observed dislocation densities correspond to those in conventional disordered alloys (Rudy, 1986; Sauthoff, 1991a). This close similarity is surprising since the elementary atomic diffusion processes and the types and numbers of the dislocation slip systems are quite different because of the different crystal symmetries. Besides dislocation creep, grain boundary sliding was observed in (Ni,Fe)Al. The process is, of course, not an independent deformation mechanism since the resulting grain shifts lead to stress concentrations at grain boundary junctions and must be accommodated by deformation processes within the grains, i.e. by dislocation creep of the grains in the stress-temperature range concerned. For such coupled deformation processes, the total creep rate is controlled by the slower process, which is dislocation creep for not too small grain sizes. With decreasing grain size the contribution of grain boundary sliding increases, from which a decrease in the total creep resistance results, as is well known for disordered alloys (Frost and Ashby, 1982) and has also been observed for (Ni,Fe)Al (Rudy, 1986; Sauthoff, 1990b). The contribution of grain boundary sliding to creep in binary NiAl has recently been studied in more detail (Raj and Farmer, 1993). At lower stresses, which produce secondary creep rates below 10 ~8 s" 1 , the observed stress-strain rate relationship deviates from power law behavior, i.e. the apparent stress exponent is smaller than 3 and decreases with decreasing stress. This deviation indicates the contribution of diffusion creep, which is a linear function of stress

706

11 Intermetallics

where Adiff is a dimensionless factor (usually Adm = 14), Q is the atomic volume, D is the effective diffusion coefficient which considers diffusion both through the grain (Nabarro-Herring creep) and along the grain boundaries (Coble creep), and d is the effective diffusion length which is usually approximated by the grain size (Frost and Ashby, 1982). Dislocation creep and diffusion creep are independent creep processes that act in parallel in the grains, and the total creep rate is given by the sum of the partial rates. Because of the stronger stress dependence of dislocation creep [Eq. (11-2)], the contribution of diffusion creep becomes more dominant with decreasing stress. In view of the observed creep behavior of (Ni,Fe)Al and other phases (Polvani etal., 1976; Nicholls and Rawlings, 1977; Ashby et al., 1978; Hirsch, 1985; Schneibel etal., 1986; Kampe etal., 1991; Hayes, 1991; Hayashi etal., 1991a; Miura etal., 1991; Rowe etal., 1991; Takahashi and Oikawa, 1991; Nathal, 1992; Nabarro and de Villiers-Filmer, 1993), it is concluded that the creep of intermetallic phases is controlled by the same creep mechanisms and is described by the same constitutive 50

equations as the creep of the familiar disordered alloys, though the crystal structures of the intermetallics are much less symmetric than the simple structures of the disordered alloys. The lower crystal symmetry leads to particular dislocation types and higher dislocation energies, including special superlattice dislocations, and to a reduced number of slip systems. However, dislocation creep is controlled by nonconservative dislocation motions, i.e. dislocation motions are coupled with local diffusion fluxes, and the latter are rate controlling. Thus the particular dislocation mechanisms are not of primary importance for the total macroscopic creep rate which explains the similar behavior of metallic and intermetallic alloys. Figure 11-23 shows the effect of composition variations, i.e. deviations from stoichiometry and variations in the Ni:Fe ratio, on the creep resistance of (Ni,Fe)Al. It can be seen that the creep resistance decreases with decreasing Al content, i.e. increasing deviation from stoichiometry. As already noted in the preceding section, deviations from stoichiometry introduce constitutional disorder, i.e. point defects which increase the diffusion coefficient

e=10'V 1 T =900°C

40-

8

30

0)

20-

c o

Figure 11-23. Creep resistance of binary and ternary B2 aluminides at 900°C (in compression with 10" 7 s" 1 strain rate) as a function of Al content (Rudy, 1986; Rudy and Sauthoff, 1986).

Q.

2? o

10 H

20

30

40

50

aluminium content in a t %

60

11.4 Nickel Aluminides and Related Phases

and thereby the creep rate. However, at lower temperatures these point defects become immobile and act as obstacles to dislocation movements, i.e. they increase the low-temperature strength. At intermediate temperatures the two effects balance each other and deviations from stoichiometry do not affect the creep resistance, as was found for NiAl (Vandervoort et al., 1966). In Fig. 11-23 the Ni-rich phases show a strong stoichiometry effect with decreasing Al content, whereas the effect is significantly smaller for the Fe-rich phases, i.e. the test temperature of 900 °C leads to high-temperature behavior by the Ni-rich phases and to intermediate-temperature behavior by the Fe-rich phases. This behavior is surprising since it is commonly supposed that such thermally activated effects scale with the melting temperature, which is higher for NiAl than for FeAl. The reasons for the observed behavior of (Ni,Fe)Al at intermediate temperatures are still to be clarified. Apart from the effect of deviations from stoichiometry, Fig. 11-23 shows a marked dependence of the creep resistance for the ternary B2 phases (Ni,Fe)Al on the Ni:Fe ratio. It can be shown for (Ni,Fe)Al that this composition dependence results from the composition dependence of the diffusion coefficient, which can be estimated by using data from Cheng and Dayananda (1979) and Moyer and Dayananda (1976) (see Rudy, 1986; Rudy and Sauthoff, 1986). The maximum in the creep resistance of (Ni,Fe)Al is directly related to the minimum in the diffusion coefficient. It has to be noted that the analysis of creep data with respect to diffusion coefficients poses problems because, apart from the scarcity of data, the diffusion coefficients in Eqs. (11-2) and (11-3) are effective ones depending on the particular creep process which determines the coupling of partial

707

diffusion fluxes, i.e. the needed effective diffusion coefficients are not those which are measured in diffusion experiments. For binary solid solutions, expressions for the effective diffusion coefficients are available for various creep mechanisms (Chin etal., 1977; Fuentes-Samaniego and Nix, 1981; Dominguez-Rodriguez and Castaing, 1993), whereas for multinary phases an expression for the effective diffusion coefficient is known only for the case of diffusion creep (Herring, 1950). Here more theoretical and experimental work is necessary. Figure 11-24 shows the correlation between the observed creep resistance of various binary and ternary B2 phases and the respective diffusion coefficients, as found in the literature or estimated on the basis of literature data (Rudy, 1986; Rudy and Sauthoff, 1986; Sauthoff, 1991 a). It can be seen that the correlation is surprisingly good and corresponds to Eq. (11-2), in

r=9oo° c -1

.E

e = 10"7s

100\ CoAl •

50-

\


Fe

D

20-

o. 2 ' A I

NiAl \

\

(N. Fe

)A(

10\

5-19

10

-17

10

FeAl \ -15

10

diffusion coefficient in m7s

Figure 11-24. Creep resistance of binary and ternary stoichiometric B2 aluminides at 900 °C (in compression with 10~ 7 s~ 1 strain rate) as a function of the diffusion coefficient, as found in the literature or estimated from available data (Rudy, 1986; Jung et al., 1987; Sauthoff, 1991a; Shankar and Seigle, 1978; Akuezue and Whittle, 1983; Cheng and Dayananda, 1979; Moyer and Dayananda, 1976).

708

11 Intermetallics

spite of the problems with the appropriate determination of the effective diffusion coefficients. The question now is why the diffusion coefficient depends in this way on the composition of the B2 phases. As for other properties, the diffusion coefficient depends on the crystal properties, i.e. on the character and strength of the atomic bonding, and indeed the activation energy of diffusion increases with increasing total enthalpy of phase formation, as is illustrated by Fig. 11-7. The question is whether such a correlation can be established - at least for a group of closely related phases - in such a way that it can be used for predicting the effects of alloying additions on the diffusion coefficients. For this the correlation of the diffusion coefficients with the character and strength of bonding must be studied in more detail. Here again much more theoretical and experimental work is necessary. In view of the discussed composition dependence of the creep resistance, it is concluded that the effective diffusion coefficient is of primary importance for controlling the creep resistance (Sauthoff, 1993 b). This of course does not mean that the other parameters in Eq. (11-2) can be neglected. This is demonstrated by the temperature dependence of the creep of B2 (Ni,Fe)Al, as was discussed earlier (Sauthoff, 1991a). In view of Eqs. (11-2) and (11-3), the apparent activation energy for creep is expected to correspond to that for diffusion since the other parameters depend less sensitively on temperature, and indeed this has been confirmed repeatedly in the case of conventional disordered alloys. However, in the case of B2 (Ni,Fe)Al, the apparent activation energy for creep only corresponds to that for diffusion at temperatures up to 900 °C, whereas at higher temperatures the apparent activation energy for creep is much higher. Acti-

vation energies for creep that are higher than those for diffusion have also been reported for other cases (Vandervoort et al., 1966; Stoloff, 1984; Whittenberger, 1986). This means that in contrast to conventional disordered alloys, the microstructuredependent parameter A in Eq. (11-2) may also exhibit a strong temperature dependence, since the shear modulus G still depends only weakly on the temperature (Harmouche and Wolfenden, 1987). The material parameter of secondary importance for creep is the shear modulus in Eq. (11-2). Elastic moduli can be calculated by quantum-mechanical, ab initio calculations (see, e.g. Fu and Yoo, 1992 a), and thus there is a physical understanding of the relation between the strength and character of bonding and the elastic behavior. However, experimental and/or theoretical data are scarce and there is little knowledge on the effects of changes in composition and structure on the elastic moduli. Figure 11-6 shows some Young's modulus data, e.g. for some B2 phases in comparison to some C15 Laves phases, which are discussed in Sec. 11.8. Obviously there is only a correlation between the Young's modulus and the total enthalpy of formation for the Laves phases. This correlation is less apparent for the B2 phases, which may be due to experimental errors and to the complexities of the elastic behavior of the B2 phases. The question is again whether such correlations can be established - at least for a group of closely related phases - and whether these correlations can then be used for predicting the effects of alloying additions on the elastic behavior. In summary, it is concluded that the creep behavior of intermetallic phases can be described as that of the conventional disordered alloys given by the familiar phenomenological constitutive equations,

709

11.4 Nickel Aluminides and Related Phases

and indeed deformation maps have been calculated as a function of available data (Jung et al., 1987). However, the physical mechanisms which control the creep behavior are only partially understood and much more work is necessary. With respect to applications as structural materials at high temperatures, the creep resistance of candidate phases should be as high as possible for increasing the service temperatures as much as possible. In view of Eqs. (11-2) and (11-3), the decisive materials' parameters for increasing the creep resistance are the diffusion coefficient D and the shear modulus G. Both parameters are determined by the atomic bonding energy to which the macroscopic enthalpy of phase formation is related (Engell et al., 1991). The latter is related to the melting temperature, and indeed there is a correlation between the parameters D and G and the melting temperature (Frost and Ashby, 1982). This means that in principle alloys with increased melting temperatures should be used for obtaining increased high temperature strengths. However, a strict validity of this rule cannot be expected since the above correlations are complex ones which depend on many parameters. An example is given by the phases NiAl and CoAl which have the same crystal structure and nearly the same melting temperature, but different creep resistances (Fig. 11-24). Diffusion depends sensitively on the crystal structure (Wever et al., 1989), and it may be supposed intuitively that a lower diffusion coefficient and thereby a higher creep resistance may be obtained by lowering the crystal symmetry. Atomic ordering in the B2 lattice results in the L2X structure with four sublattices, which is the case for the Heusler-type phases Ni2TiAl and Co2TiAl. Indeed Co2TiAl shows a much higher flow stress than (Co 0 8 Fe 0 2)A1 according to data in Sauthoff (1989), and

Ni2TiAl is more creep resistant than NiAl (Strutt and Polvani, 1973). Besides crystal symmetry, atomic packing should also be considered, and it is well known that, e.g. the self-diffusion coefficient of Fe in the close-placked f.c.c. lattice is lower by two orders of magnitude than that in the more open b.c.c. structure (Fridberg et al., 1969). The B2 structure results from ordering in the b.c.c. lattice, whereas ordering in the f.c.c. lattice leads, e.g. to the L l 2 structure of Ni 3 Al or with carbon to the L'l 2 structure of Fe 3 AlC (Jung and Sauthoff, 1989 b) and, furthermore, ordering in the L l 2 lattice leads to the D0 22 structure of Al 3 Nb, as is shown in Fig. 11-1. Figure 11-25 combines creep data for all these structures, and it can be seen that the differences in crystal structure with respect temperature in °C 700 800 9001000 1200 1400 Ni3Al

LJ

Al3Nb

100

c

a to

a.

10 £ = 10 S

1.0

0.8

0.6

1000/7" in K" Figure 11-25. Temperature dependence of creep resistance (in compression with 10" 7 s~1 secondary strain rate) for various single-phase intermetallic alloys: NiAl, CoAl and related alloys with a B2 structure (Jung etal., 1987; Sauthoff, 1989), Ni2TiAl with an L2X structure (Strutt and Polvani, 1973), two Ni3Al variants with Ll 2 structures, i.e., an advanced aluminide ( + ) (Schneibel et al., 1986) and Ni 3 Al-Fe (x) (Nicholls and Rawlings, 1977), Fe3AlC with an L'l 2 structure (Jung and Sauthoff, 1989 b), and Al3Nb with a D0 22 structure (Sauthoff, 1990a, b; Reip, 1991).

710

11 Intermetallics

to symmetry and atomic packing indeed result in correspondingly significant differences in the creep resistance. However, it can also be seen that the variation in the creep resistance with varying sublattice occupation in the B2 structure is as large as the variation with varying lattice structure, i.e. both effects have to be considered for increasing the high temperature strength. Creep Resistance of Multiphase NiAl-Base Intermetallic Alloys Precipitated particles are usually used in conventional alloys for increasing the creep resistance. The strengthening effect of precipitated particles was studied in NiAl-Fe alloys with a B2 NiAl matrix and a-Fe particles (Jung and Sauthoff, 1989 a). It was found that the effect of the a-Fe particles was quite analogous to that of B2 NiAl particles in Fe-NiAl alloys with an a-Fe matrix, which were studied by Jung and Sauthoff (1987) and Jung (1986). In both cases the particles act as dislocation obstacles because of the dislocation-particle interaction, and indeed adhering dislocations were observed at particles before detachment in the NiAl-Fe case. Such obstacles are surmounted by climb which gives rise to a threshold stress ath and increases the creep resistance according to (11-4) kT This threshold stress is proportional to the Orowan stress, as was shown theoretically for various climb processes (Arzt and Rosier, 1988), and thus is proportional to the reciprocal particle distance, in agreement with the experiments (Jung and Sauthoff, 1987,1989 a). The observed, secondary dislocation creep can be described by Eq. (11-4) and deformation maps can be calculated on the basis of the experimental data (Jung and Sauthoff, 1989 a).

In particulate alloys one phase, i.e. the matrix, is distributed continuously, whereas the second phase is distributed discontinuously. In nonparticulate alloys all phases are distributed continuously, as is the case in lamellar or fibrous composites. The effect of phase distribution on the creep behavior of NiAl-Fe alloys with lamellar microstructures was studied in detail (Klower, 1989; Klower and Sauthoff, 1991, 1992). For this an Ni-40 at.% Fe-18 at.% Al alloy was chosen which was produced by directional solidification to give a lamellar microstructure of the phases B2 NiAl and disordered f.c.c. y-Fe-Ni with equal volume fractions. It was found that the creep resistance of such a lamellar alloy is related to the creep resistance of the constituent phases according to a rule of mixtures as long as the lamellae spacing is larger than a critical spacing, which is of the order of the free dislocation path. If the lamellae spacings are smaller than this critical value, then the lamellae interfaces give rise to an additional strengthening effect. This strengthening effect can be described, as in Eq. (11-4), by a threshold stress which again is proportional to the reciprocal lamellar spacing. Similar effects have been observed in other intermetallic NiAl-base alloys with less regular distributions of hexagonal C14 Laves phases (Machon, 1992; Sauthoff, 1993 a), and have been discussed by Sauthoff (1991b). In those alloys with coarse phase distributions the observed secondary creep rates follow a rule of mixtures at a first approximation, and additional strengthening effects are only observed for alloys with fine phase distributions. From this it is concluded that particulate and nonparticulate intermetallic alloys creep in similar ways and can be described by the same constitutive equations as conventional multiphase alloys.

11.4 Nickel Aluminides and Related Phases

However, the above discussion has only referred to secondary creep. The situation is less clear for the transient primary creep stage which precedes secondary creep. As already discussed (Sauthoff, 1991b), the primary creep strain is reduced significantly by the presence of second phases and it decreases with increasing stress and with decreasing interface spacing, at least in some NiAl-base alloys (Reip, 1991; Klower, 1989). Such a behavior is also known for conventional alloys, but there are other alloys which show an opposite behavior, i.e. an increasing primary strain with increasing stress (see Sauthoff, 1991b). Such effects are not yet understood even for disordered alloys. It is further noted that not only the normal primary creep with a decelerating creep rate is observed, but also inverse primary creep with an accelerating creep rate, e.g. in the case of ternary Laves phases (Sauthoff, 1990a, 1991a; Machon, 1992). Such inverse primary creep has also been reported for Ni 3 Al, where it is due to the activation of various slip systems (Hazzledine and Schneibel, 1989; Schneibel and Hazzledine, 1992). Inverse creep results from an insufficient number of mobile dislocations, and a classic example for such a deformation behavior is silicon, where the deformation behavior has been analyzed in detail (Alexander and Haasen, 1968; Alexander, 1986). In any case, the creep resistance of NiAl can be increased significantly by alloying with other stronger intermetallics to form multiphase intermetallic alloys which may be regarded as intermetallic in situ composites. This is exemplified by Fig. 11-26 with creep resistance data for various N i Al-Nb alloys which contain the tetragonal D0 22 phase Al 3 Nb and the hexagonal Laves phase NbNiAl with a C14 structure, as well as NiAl. It can be seen that the

711

temperature in °C 800 900 1000 1100

10-

1.0

0.9

0.8

0.7

reciprocal temperature in 10" K"

Figure 11-26. Creep resistance (in compression with 10~ 7 s~ 1 secondary strain rate) as a function of temperature for various NiAl-Al 3 Nb alloys, i.e. for the three-phase alloys NiAl-Nb-1, -2, and -3 with 22 vol.%, 46 vol.%, and 78 vol.% NiAl, and 4 vol.%, 15vol.%, and 15 vol.% NbNiAl, respectively, compared with single-phase NiAl (Rudy and Sauthoff, 1985; Rudy, 1986), and Al3Nb and Al 3 (Nb 075 Ti 0 25) - both with a D0 22 structure (Reip, 1991).

creep resistance of these intermetallic alloys increases continuously with increasing volume fractions of stronger phases from that of the comparatively soft NiAl to that of the hard phases Al 3 Nb and Al3(Nb,Ti), which have been discussed in Sec. 11.3.3.2 [the Al3(Nb,Ti) data have been introduced into this figure because only this ternary trialuminide could also be tested at lower temperatures]. Another example is given in Fig. 11-27, which illustrates the strengthening of NiAl by ternary C14 Laves phases. The data points for TaNiAl-NiAl seem to follow a rule of mixtures, but the data for NbNiAl NiAl clearly indicate a significant deviation from a lineal superposition for the creep resistances. Furthermore, the extrap-

712

11 Intermetallics

T = 1200°C E --io-V

30-

NbNiAl P

1

TaNiAl-NiAl

QJ

Figure 11-27. Creep resistance (compressive stress for 10~7 s" 1 secondary creep rate) at 1200 °C as a function of the volume fraction of the second phase for the intermetallic NiAl-base alloys NiAlNbNiAl and NiAl-TaNiAl with the C14 Laves phases NbNiAl and TaNiAl, respectively (Sauthoff, 1990 a; Machon, 1992).

20

|

NbNiAl-NiAl QJ

(D QJ

10 • NiAl

0

0 NiAl

25 50 75 volume fraction in %

olated creep resistance for 0 % strengthening phase is appreciably higher in this case than the measured value for binary NiAl. This additional strengthening effect is due to precipitated, fine particles of the Heuslertype phase Nb 2 NiAl with an L2X structure, which give rise to a threshold stress (Sauthoff, 1991b; Machon, 1992). Further examples are given by Nathal (1992), where a wealth of data from various alloy systems based on the B2 phases are analyzed. 11.4.3.5 Environmental Effects

Unlike most other aluminides which are considered for high temperature applications, NiAl with a B2 structure exhibits excellent oxidation resistance since a protective A12O3 scale is readily formed during oxidation (Doychak et al., 1989; Nesbitt and Lowell, 1993). According to present knowledge, NiAl seems to be the only really oxidation resistant intermetallic, apart from some silicides (Meier et al., 1993). The physical reason for this high oxidation resistance is that the Al content

100 NbNiAl TaNiAl is sufficiently high and the Al diffusion is sufficiently fast in NiAl at all temperatures to form stable A12O3 scales at the surface and avoid internal oxidation in the bulk (Pettit, 1967). The details of scale growth have been described by Doychak et al., (1989). During oxidation voids are formed at or near the NiAl/Al 2 O 3 interface because of unbalanced diffusion fluxes, corresponding to the Kirkendall effect (Cathcart, 1985; Meier et al., 1993). This void formation, which has also been observed for other Al-rich alloys, leads to spalling of the A12O3 scale and affects the oxidation resistance. Void formation is enhanced by the segregation of sulfur to the free NiAl surface of the voids, thus giving rise of the deleterious "sulfur effect" (Grabke et al., 1990, 1991b). Scale adherence is also affected by the presence of TiB 2 dispersoids, as has been shown for an NiAl-TiB 2 composite (Pregger et al., 1992). As in many other cases, scale adherence and oxidation resistance are improved by microalloying with "oxygen active" elements, e.g. Y, Zr, Hf, Ce, or La (Mrowec

11.4 Nickel Aluminides and Related Phases

and Jedlinski, 1989; Grabke et al., 1991 b). In particular, Zr has been used to increase the resistance of NiAl to cyclic oxidation (Doychak et al., 1989; Nesbitt et al., 1992). The physical understanding of the beneficial effect of such "oxygen active" elements is still insufficient, as is also the case for conventional disordered alloys. The effect of alloying with Cr and Si, which are used in other alloys for providing oxidation resistance, has been reviewed (Meier, 1989). It is noted that the sulfidation resistance of NiAl is reduced by Cr (Mrowec etal., 1989). Apart from oxidation with external scale formation, NiAl, as well as various other intermetallics, is subject to a special oxidation phenomenon which is known as pesting (Aitken, 1967; Meier and Pettit, 1992). Pesting occurs in a critical temperature range - usually at about 800 °C - and may lead to complete disintegration of the material. NiAl is less susceptible to pesting than other intermetallics, and it shows pesting only at very low oxygen pressures when oxygen penetrates the thin surface scale and diffuses inwards along the grain boundaries (Grabke et al., 1991 a, b; Meier etal., 1993). Environmental embrittlement, which affects the mechanical behavior of various other intermetallics, e.g. the other B2 aluminide, FeAl, has not been found for NiAl (Liu, 1992; Lahrman et al., 1993 b). 11.4.3.6 Alloy Developments and Applications

Magnetic Alloys The advantageous magnetic properties of Fe-Ni-Al alloys with a suitable composition were first discovered by Mishima in 1931, which initiated the development of the Alnico alloys for applications as permanent magnet materials (see, e.g. Jelling-

713

haus, 1936,1943; De Vos, 1969). The basic alloy contains about 50 at.% Fe, 25 at.% Ni, and 25 at.% Al to which various other elements - Co or Cu as macroalloying elements and Nb or Ti as microalloying elements - are added for property optimization (McCurrie, 1986). The alloys are single-phase at high temperatures and decompose spinodally after cooling at a critical rate to form a two-phase microstructure with an interlocking phase distribution. The constituent phases of these alloys are an NiAl-base phase with a B2 structure and a disordered b.c.c. Fe-rich phase, according to the Ni-Al-Fe phase diagram (Bradley and Taylor, 1938; Dannohl, 1942; Bradley, 1951; Hao etal., 1984). The B2 phase in the Alnicos is only weakly ferromagnetic whereas the Fe-rich phase is strongly ferromagnetic, and thus the interlocking phase distribution and the strong shape anisotropy of the Fe-rich phase results in a high coercivity (De Vos, 1969; McCurrie, 1986). The magnetic behavior is a sensitive function of the pretreatment, not only with respect to the phase distribution, but also with respect to the concentration and distribution of the defects, e.g., vacancies and antistructure atoms (Kilner and Harris, 1981). The effect of the additional alloying elements Co, Mn, Ti, or Cu which dissolve preferentially in the B2 phase and Cr, Mo, V, Nb, or Si which dissolve preferentially in the disordered Fe-rich phase - on the decomposition process of the Alnico alloys was studied by Hao et al. (1985). Alnico alloys are brittle and hard, they can only be machined by grinding, spark erosion, and electrochemical milling, and they resist atmospheric corrosion well up to 500 °C (Fiepke, 1990). The mechanical behavior, in particular creep, of Alnicotype, Fe-Ni-Al alloys has been studied in detail, and both Fe-rich alloys with precip-

714

11 Intermetallics

itated NiAl particles and NiAl-rich alloys with precipitated Fe particles have been considered (Jung and Sauthoff, 1987, 1989 a). The magnetic properties are summarized and compared to those of other alternative alloys by Fiepke (1990) and McCurrie (1986). Alnico alloys are superior to other permanent magnet materials at resisting temperature effects on magnetic properties. It is noted that presently NiAl is also considered an attractive candidate phase for other functional applications, e.g. contacts in electronic thin film devices or high voltage electrodes in electron-optical instruments (Miracle, 1993). Shape Memory Alloys It is well known that the martensitic transformation of Al-deficient NiAl (see Sec. 11.4.3.2) is thermoelastic and produces the shape memory effect. Consequently materials developments have been started which aim at applications as shape memory alloys (Furukawa etal., 1988; Kainuma et al., 1992b, c). The martensitic transformation temperature can be varied within a broad temperature range up to 900 °C, and thus the shape memory effect can be produced at high temperatures which allows the development of hightemperature shape memory alloys. The problem of low room temperature ductility of NiAl has been overcome by alloying with a third element - in particular Fe - to produce a ductile second phase with an f.c.c. structure. Coating Alloys Many high-temperature alloy components with insufficient resistance to corrosion in hot gases, e.g. in gas turbines, are coated with a corrosion resistant surface

layer for providing the necessary protection from the particular corrosion (see, e.g. Griinling et al., 1983; Pettit and Goward, 1983; Nicholls and Stephenson, 1991). In view of the exceptionally high oxidation resistance of the B2 phase NiAl (see Sec. 11.4.3.5) it has long since been used as a coating material. Indeed, one of the oldest processes for producing coatings is pack aluminizing, by which Al diffuses from the Al-rich pack into the alloy to be coated and forms an aluminide surface layer, which consists of NiAl in the case of Ni base alloys (Patnaik, 1989). Such diffusion coatings are still extensively used for protecting turbine blades under various turbine conditions. The coating performance is improved by adding alloying additions, i.e Cr, Si, Ta, rare earths, or precious metals, which is achieved by modifying the pack aluminizing process (Nicholls and Stephenson, 1991). Besides such diffusion coatings, overlay coatings are used which are produced by depositing a specifically designed corrosion resistant alloy onto the component surface (Nicholls and Stephenson, 1991). These coating alloys, which are described by the general composition formula MCrAlY with M = Fe, Ni, and/or Co again rely on the presence of MAI phases with a B2 structure (see, e.g. Gudmundsson and Jacobson, 1988). The advantage of these coatings is the negligibly small chemical interaction between the coating and the substrate during deposition, the choice of the corrosion resistant alloy composition and the ability to deposit thicker coatings for extended service lives. It is noted that during service interdiffusion occurs between the coating and the substrate from which an Al depletion of the coating may result. Al deficient NiAl may transform martensitically because of cooling cycles, which promotes spalling,

11.4 Nickel Aluminides and Related Phases

i.e., it contributes to the degradation of the coating (Smialek and Hehemann, 1973). Structural Alloys For a long time NiAl has been considered as a basis for developing structural materials for high temperature applications (Fitzer and Gerasimoff, 1959; Imai and Kumazawa, 1959; Grala, 1960; Jellinghaus, 1961, 1967). NiAl is advantageous because of its low density, good thermal conductivity, high melting point, and excellent oxidation resistance, whereas its little deformability at room temperature and its low strength and creep resistance at high temperatures above 1000 °C is always a disadvantage compared to the superalloys (Vedula and Stephens, 1987 a; Darolia, 1991). The low strength at high temperatures is related to the excellent oxidation resistance, since both effects result from easy diffusion at high temperatures which is due to the open B2 structure of NiAl (see Sees. 11.4.3.4 and 11.4.3.5). In view of applications as high-pressure turbine blade material in flying gas turbines, NiAl is regarded as highly promising for developing alloys which are competition for superalloys because the strength deficit is compensated for by the advantage in density and thermal conductivity, and respective materials developments are under way (Darolia etal., 1992b; Liu and Kumar, 1993; Walston and Darolia, 1993; Locci et al., 1993; Field et al., 1993; Goldman, 1993; Igarashi and Senba, 1993; Oti and Yu, 1993; Darolia, 1993). Microalloying additions of Mo, Fe, Nb, Ta, Zr, Hf, B, and C as well as special processing routes, e.g. hot extrusion after melting and casting, are used, and the aim is always improved room temperature ductility and increased high temperature strength (Walston etal., 1993; Field etal., 1993; Lahrman etal., 1993a;

715

Locci etal., 1993; Hack etal., 1993). The mechanical behavior of NiAl under complex conditions has not yet been studied systematically, as is exemplified by the only work on fatigue to date (Noebe and Lerch, 1992, 1993; Smith etal., 1992; Stoloff, 1992; Cullers et al., 1993; Edwards and Gibala, 1993). As in the case of conventional disordered alloys, NiAl alloys can be strengthened appreciably by second phases. Apart from the effects on strength, second phases may also be beneficial for ductility and toughness, as has been discussed with respect to NiAl-based intermetallic alloys (Noebe et al., 1991; Clemens and Bildstein, 1992). In any case, the effects of second phases depend on the properties of the respective phases and on the phase distribution. This is illustrated in the following sections by regarding various NiAl alloy systems which have been studied in some more detail. The creep behavior of NiAl alloys with strengthening second phases is addressed in Sec. 11.4.3.4. NiAl-Fe Alloys: The ternary Ni-Fe-Al phase diagram in Fig. 11-28 shows various two-phase and three-phase fields neighboring the (Ni,Fe)Al field, i.e. there are various possibilities for producing NiAl alloys with second phases. For example, fine b.c.c. Ferich particles can be precipitated from the NiAl-Fe matrix by careful adjustment of the Fe and Al contents and by adequate heat treatment (Jung and Sauthoff, 1989 a). The strengthening effect of these disordered particles in the ordered intermetallic matrix is quite analogous to that of ordered NiAl precipitate particles in a disordered Fe-base alloy (Jung and Sauthoff, 1987), and thus corresponds to particle strengthening in conventional alloys. The

716

11 Intermetallics

80

60

80

'100 100 A l

composition in at.% Figure 11-28. Isothermal section of the Ni-Fe-Al phase diagram at 400 °C with the f.c.c. Ni solid solution (Al), the b.c.c. oc-Fe solid solution (A2) and the intermetallic phases Ni3Al (Ll 2 ), Fe3Al (D03) and (Ni,Fe)Al (B2 or L20) [schematic diagram from Sauthoff (1986) according to Bradley and Taylor (1938), Dannohl (1942), Bradley (1951), and Hao et al. (1984)].

two-phase equilibrium between NiAl and disordered f.c.c. y-Fe-Ni has been used to produce lamellar two-phase alloys, i.e. in situ composites, by directional solidification (Klower and Sauthoff, 1991). The strength is described by a rule of mixtures as long as the lamellar spacing is large, i.e. for coarse microstructures, whereas an additional strength increase is obtained for fine microstructures with small lamellar spacings (see Sec. 11.4.3.4). Besides strengthening, second phases may also be useful for improving the ductility and the toughness - in particular if they are soft, as is the case for the abovementioned, disordered phases (Noebe e t a l , 1991). Even the two-phase eqilibrium between NiAl and the other nickel aluminide Ni 3 Al can be used to produce a two-phase NiAl-Ni 3 Al alloy with improved strength and toughness (Baker and

Munroe, 1990). Consequently, the composition range of the Ni-Fe-Al system, which includes the phase equilibria between NiAl, a-Fe, y-Fe and Ni 3 Al, has been studied intensively in order to identify NiAl-base alloys with good combinations of strength and ductility, and promising materials developments have been started (Guha et al., 1989,1992; Baker and George, 1992; M. Larsen et al., 1990; Noebe et al., 1991; Raj et al., 1991,1992b; Raj, 1992; J. H. Lee etal., 1992; Tsau etal., 1992; Golberg and Shevakin, 1991; Misra etal., 1993; Kostrubanic etal., 1993). In particular, alloys with 20 at.% Al and 30 at.% Fe, as well as with 30 at.% Al and 20 at.% Fe, are subjects of materials developments. The fatigue behavior has been studied by Stoloff (1992). It has to be emphasized that both composition and processing have to be controlled carefully, since the balance of strength and ductility depends sensitively on the distribution of the phases. It is noted that similar Ni-Co-Al alloys have been prepared which also offer good possibilities for optimizing the strength and ductility (Kimura et al., 1993). The NiAl phase in NiAl alloys with second Fe-rich and/or Ni-rich phases is offstoichiometric with an Al deficiency. Thus NiAl in these alloys may transform martensitically, as has been discussed in Sec. 11.4.3.2, which offers a further possibility for improving the strength and ductility. Indeed such alloys with martensitically transformable NiAl exhibit some room temperature tensile ductility (Khadikar etal., 1987; Furukawa etal., 1988; Ishida etal., 1991a; Kainuma etal., 1992 b). It is noted that these alloys show a shape memory effect and are of interest in view of respective applications.

11.4 Nickel Aluminides and Related Phases

NiAl-Cr Alloys: As has been already discussed (Sauthoff, 1990 a), NiAl forms eutectics with the refractory metals Cr, Mo, and W which may be used to produce composite materials by directional solidification (Cline et al., 1971), and indeed the NiAl-Cr composite shows promising high-temperature strength (Walter and Cline, 1970). Such alloys - in particular NiAl-Cr and NiAl-Mo - are now studied intensively in order to develop high-temperature, high-strength alloys (Kowalski and Frommeyer, 1992; Subramanian et al., 1990a; D. R. Johnson et al., 1992, 1993; Bowman, 1992; Chang, 1992; Heredia and Valencia, 1992; Goldman, 1993). For the preparation of alloys, both powder metallurgy and ingot metallurgy in particular directional solidification - are used. Compressive and tensile strength and toughness have been studied and preliminary data have been presented in Sauthoff (1990a). Strength increases significantly with increasing Cr content, i.e. increasing volume fraction of the disordered b.c.c. aCr phase in these two-phase alloys. Furthermore, strength - in particular at high temperatures - increases with increasing grain size, whereas ductility decreases. In any case ductility and toughness are sensitive functions of the microstructure and can be improved by appropriate thermomechanical processing. This means on the one hand that the processing of such alloys must be controlled carefully for a good and reproducible materials' quality, and on the other hand there are good chances of increasing the low toughness by optimizing the alloy with respect to composition, phase distribution and microstructure. The fracture resistance of directionally solidified NiAl-Cr and NiAlMo alloys has recently been studied (Here-

717

dia et al., 1993). It has been found that the crack-initiation toughness is appreciably higher than that of NiAl, which has been rationalized by available models. The excellent oxidation resistance of pure NiAl decreases with increasing Cr content and increases with increasing temperature (Brumm and Grabke, 1992; Grabke et al., 1992). In any case the oxidation resistance of NiAl-Cr alloys is still acceptable for not too high volume fractions of the oc-Cr phase. NiAl-Ti Alloys: Alloying of NiAl with Ti produces the hard and brittle Heusler-type phase Ni2AlTi with an L2X structure (Villars and Calvert, 1991). It is very stable and strongly bonded which makes deformation difficult, and the strength increases with deviations from stoichiometry (Umakoshi etal., 1985). Ni2AlTi forms a stable twophase equilibrium with NiAl (Raman and Schubert, 1965; Nash and Liang, 1985; Mazdiyasni et al., 1989; Kumar, 1990; Lee and Nash, 1991b) and accordingly twophase alloys can be produced with either Ni2AlTi precipitate particles in an NiAl matrix or NiAl precipitate particles in an Ni2AlTi matrix. Such alloys were studied intensively with respect to creep (Strutt and Polvani, 1973; Polvani et al., 1976). In particular it was found that the creep resistance of alloys with an Ni 2 AITi matrix and small NiAl particles reaches that of the Ni-base superalloy MARM-200. The microstructure of the Ni 2 AlTiNiAl alloys is quite analogous to that of superalloys with semicoherent phase boundaries and analogous particle shapes and orientation relationships. This semicoherency and the orientation relationship result from the fact that the B2 structure and the L2X structure are closely related

718

11 Intermetallics

ably to superalloys with respect to strength. The slip transfer from the NiAl phase to the Ni3Al phase in such alloys has been studied in detail (R. Yang et al., 1993). In view of these properties and the comparatively low density in the range 6.6-6.9 g/cm3, the three-phase NiAlNi 2 AlTi-Ni 3 Al alloys are regarded as promising for applications in aero-engines (YangetaL, 1992 a).

(see Fig. 11-1), and the lattice misfit between the two phases is small (Takeyama et al, 1991). However, the lattice misfit is not negligible since it gives rise to an elastic energy contribution which controls the microstructure evolution (Bendersky et al., 1988). The oxidation rate of Ni2AlTi is of the same order of magnitude as that of dilute y-Ni-Al alloys (Lee and Shen, 1989). Presently, a materials development on the basis of NiAl with strengthening Ni2AlTi is under way which aims at applications such as turbine blades in aeroengines (Darolia, 1991; 1993) (see Sec. 11.4.3.6.). The problem with such alloys is their brittleness. Recently it has been found that the brittleness can be relieved by making use of the three-phase equilibrium between NiAl, Ni2AlTi, and Ni 3 Al (Yang et al., 1992 a, b). The resulting three-phase alloys are plastically deformable in compression at room temperature and compare favor-

NiAl-Nb and NiAl-Ta Alloys: In view of the ternary N i - A l - N b phase diagram (see Fig. 11-29) there are various possible ways to form two-phase or threephase alloys on the basis of NiAl by alloying with Nb. For the strengthening of NiAl, second phases with not too low Al contents are preferred with respect to density and oxidation resistance. Al 3 Nb with a tetragonal D0 22 structure, which has been discussed in Sec. 11.3.3.2, and the Laves phase NbNiAl with a hexagonal C14

Figure 11-29. Isothermal section of the Ni-Al-Nb phase diagram at 1140°C (Benjamin et al., 1966). 20

30

40

50

— at. % Al-

60

70

80

719

11.4 Nickel Aluminides and Related Phases

structure, which will be discussed in Sec. 11.8, are such phases, and both phases are significantly stronger and more brittle than NiAl. Figure 11-30 a shows flow stress data as a function of temperature for these phases and for three three-phase alloys which are composed of NiAl, Al 3 Nb and NbNiAl and are either NiAl-rich or Al3Nb-rich. It can be seen that NiAl is indeed strengthened significantly by the hard phases Al 3 Nb and NbNiAl. At high temperatures the flow stress increases with increasing volume fraction of the hard phases according to a rule of mixtures at a first approximation. At low temperatures an inverse behavior is observed. This may be related to the fact that the phases - in particular NiAl - which are in mutual equilibrium have off-stoichiometric compositions, and the resulting constitutional defects which are immobile at low temperatures give rise to additional strengthening. It is noted that the flow stress curves for Al 3 Nb and NbNiAl do not extend to low temperatures because at low temperatures the fracture strain is smaller than the 0.2 % proof strain. The three-phase alloys can be tested at much lower temperatures because microcracks in the hard phases are stopped at the phase boundaries of NiAl at intermediate temperatures. Indeed the apparent brittle-to-ductile transition temperature of these alloys is between that of NiAl at about 400 °C, and that of Al 3 Nb and NbNiAl at about 1100°C, as is shown in Fig. 11-30 b. The oxidation resistance of these alloys with Al 3 Nb has also been studied in some detail, and it was found that the good oxidation resistance of binary NiAl can nearly be reached by proper adjustment of the volume fractions of the respective phases (Steinhorst, 1989; Steinhorst and Grabke, 1989; Grabke etal., 1990, 1991b, 1992).

2000 Al-Nb-Ni-3 AWMb

0

200 400 600

800 1000 1200 1400

temperature

(a)

5

Al-Nb-Ni-2 t Al-Nb-Ni-3

in °C

At-Nb-Ni-1

°E 4 a -t 3 ai

£ 2a ^

1-

800

(b)

900

1000

1100

1200

1300

temperature in °C

Figure 11-30. (a) 0.2% proof stress (in compression with 10" 7 s" 1 secondary strain rate) and (b) fracture strain (in bending with 10~ 4 s~ 1 strain rate - always referred to the specimen edge) as a function of temperature for various NiAl-Al 3 Nb alloys, i.e. for the three-phase alloys NiAl-Nb-1, -2, and -3 with 22vol.%, 46vol.%, and 78 vol.% NiAl (B2), and 4vol.%, 15vol.%, and 15 vol.% NbNiAl (C14), respectively - balance'Al3Nb (D022) - compared with single-phase NiAl, Al3Nb and NbNiAl (Rudy and Sauthoff, 1985; Rudy, 1986; Sauthoff, 1990a, b; Reip, 1991; Machon, 1992; Reip and Sauthoff, 1993).

720

11 Intermetallics

In spite of the obtained strength increases with Al 3 Nb, the strength of the NiAl alloys with Al 3 Nb - in particular the creep resistance - is not yet sufficient to enter the temperature range above that of the superalloys. Thus one has to rely on the Laves phase NbNiAl, which can be precipitated from NiAl, to obtain particle-strengthened NiAl alloys (Sherman and Vedula, 1986; Vedula and Stephens, 1987b). NbNiAl is significantly stronger than Al 3 Nb, but also more brittle, and crack-free NbNiAl specimens can be produced only by powder metallurgy. However, the Laves phase forms a stable equilibrium with the B2 phase NiAl (Fig. 11-29), and the resulting NbNiAl-NiAl alloys, which have been studied in some detail (Sauthoff, 1990 a, b, 1993 a; Machon, 1992), can be produced without defects by vacuum induction melting. The brittleness of these two-phase alloys is reduced to such an extent that the alloys can be hit by a hammer at room temperature without cracking, i.e. they can be handled safely. Figure 11-31 shows yield strength data of two-phase NiAl-NbNiAl alloys as a function of temperature for increasing volume fractions of the strengthening Laves phase. It is noted that the Laves phase precipitates on the grain boundaries and

forms a continuous skeleton in all alloys so that the Laves phase plays the role of the matrix even in the NbNiAl-poor alloys. The compressive strength again follows a rule of mixtures to a first approximation. The bending strength does not increase with the NbNiAl volume fraction at higher volume fractions of NbNiAl. This is supposed to be due to alloy defects, i.e. pores and microcracks, and thus further work has to concentrate on improving the materials' quality. Up to now the fracture toughness at room temperature is 2-4 MNm~ 3/2 (see data in Sauthoff, 1990 a, b) for these alloys, and the brittle-to-ductile transition temperature is between about 500 °C and 700 °C depending on composition and preparation method. Analysis of the microstructures of the deformed alloys has shown that in the brittle-to-ductile transition range cracks in the Laves phase are stopped at the Laves phase-NiAl interface by producing a plastic zone in the NiAl phase ahead of the crack (Machon, 1992; Wunderlich et al., 1992). Recently similar alloys have been prepared by mechanical alloying (Arzt et al., 1993; Clemens etal., 1993 a). Fully dense and completely aligned, lamellar, in situ composite NiAl-NbNiAl alloys have been

1200c

1000-

CO cu c

800600-

M—

o o

c , CL

ft?

400200-

0200

400

600

800

1000

temperature in °C

1200

1400

Figure 11-31. 0.2% proof stress in compression (10 deformation rate) as a function of temperature for NiAl (B2), the Laves phase NbNiAl (C14) and various two-phase alloys with these two phases (the percentages at the curves are the respective amounts of NbNiAl) (Sauthoff, 1989, 1990 a, 1991a; Machon, 1992) compared with data for the ODS superalloy MA 6000 (in tension) (Inco, 1982).

721

11.4 Nickel Aluminides and Related Phases

produced successfully by directional solidification (Reviere etal., 1992; Whittenberger etal., 1992a). This special microstructure leads to a further increase in the strength, but has no beneficial impact on the low-temperature ductility and fracture strength. Other ternary Laves phases with Al are NbFeAl, TaFeAl, and TaNiAl (see Sec. 11.8.1), and the latter in particular has been considered for strengthening NiAl (Pathare et al., 1987; Vedula and Stephens, 1987b). These phases are still harder than NbNiAl, they also form stable equilibria with the respective B2 aluminides FeAl and NiAl, and preliminary data for the respective one-phase and two-phase alloys have been presented in Sauthoff (1990 a, b, 1993 a). Figure 11-32 shows yield strength data for two-phase NiAl-TaNiAl alloys. It can be seen that the obtained flow stresses are significantly higher than those of the NiAl-NbNiAl alloys, and thus it should be possible to enter the strength-temperature range above that of superalloys by making use of such alloys. The Arrhenius plot of the data shows that a rule of mixtures may be supposed not only for the yield strength, but also for the apparent activation energy. It has to be noted, however, that these alloys compare less favorably with superalloys with respect to longterm behavior, i.e. creep strength, which has been addressed in Sec. 11.4.3.4. Since the transition metal elements in these ternary Laves phases can substitute for each other freely, and since the distribution of Laves phase can be controlled by thermomechanical treatments there are possibilities for optimizing such NiAl alloys with Laves phases with respect to creep resistance and the brittle-to-ductile transition temperature (BDTT) by controlling the composition and phase distribution, and this is being studied presently

temperature in °C 1400

1200

1000 900

800

700

/ D

TaNiAl

/

f

10'

^ , - ^ N i A l A

10

10 i/y-10' 4 in K"1

Figure 11-32. 0.2% proof stress in compression (10" 4 s~ 1 deformation rate) as a function of temperature for NiAl (B2), the Laves phase TaNiAl (C14) and various two-phase alloys with these two phases (the percentages at the curves are the respective amounts of TaNiAl) (Sauthoff, 1989, 1990a, 1991a; Machon, 1992).

(Zeumer etal., 1991; Sauthoff, 1993a). Variation of the BDTT with changing alloy composition is illustrated in Fig. 11-33. Finally, it is noted that besides the Laves phases the Heusler-type phases Ni 2 AlNb and Ni2AlTa with an L2X structure may BDT temperature in °C 1200

1100

1000

900

\ -Nb-rich

£ 10"

-

Ta-ri ch

•

brittle

\ \

\ ^

duct le

_

-i

10" 0.6

0.7

0.8

0.9

1/7 BDTT in [ 1 0 0 0 / K ]

Figure 11-33. Brittle-to-ductile transition temperature BDTT as a function of strain rate in bending (referred to the specimen edge) for various two-phase NiAl-(Ta,Nb)NiAl alloys with 23.5 vol.% Laves phase: TalONi45A145 (x), Ta9NblNi45A145 ( + ), Ta7.5Nb2.5Ni45A145 (o), TalNb9Ni45A145 (#), and NblONi45A145 (*) (Sauthoff, 1991a; Zeumer etal., 1991; Zeumer and Sauthoff, 1992).

722

11 Intermetallics

form in the ternary systems Ni-Al-Nb and Ni-Al-Ta, respectively, in close analogy to the Ni-Al-Ti system, as is shown in the respective phase diagrams [see Fig. 11-29 and Pak et al. (1988) and Darolia (1991)]. Indeed two-phase NiAl alloys with these strengthening precipitates have been studied with respect to mechanical behavior and are regarded as promising for materials developments (Pak et al., 1988; Yasuda et al., 1992), as well as threephase NiAl alloys with a coarse distribution of the NbNi Al Laves phase and a fine particle distribution of Ni 2 AlNb (Machon, 1992; Machon and Sauthoff, 1994). In summary, it is concluded that there are good prospects for developing materials on the basis of NiAl with strengthening Laves phases and/or Heusler-type phases which have sufficient creep strength at temperatures above the application temperatures of superalloys and tolerable brittleness. ODS NiAl Alloys and Composites: In view of the advantageous properties of NiAl for high-temperature applications, i.e. excellent oxidation resistance, moderate density, high stability, and good thermal conductivity, and the disadvantageous poor strength at high temperatures, NiAl has been considered as a matrix material for intermetallic matrix composites, and respective materials developments are under way (see reviews by Bowman and Noebe, 1989; Rigney et al., 1989; Kumar, 1991; Vedula, 1991; Kumar and Whittenberger, 1992; Kumar et al., 1992b; Shah and Anton, 1993). Such composite materials have been produced primarily by powder metallurgy methods, and a variety of phases have been used for strengthening the NiAl matrix, i.e. A12O3 and other oxides (Jellinghaus, 1961; Noebe etal., 1990; Al-

man and Stoloff, 1991; Dimiduk etal., 1991; Kostrubanic etal., 1991; Kumar, 1991; Nourbakhsh et al., 1991; Baker and George, 1992; Dymek etal., 1992; Bowman, 1992; Anton and Shah, 1992 b; Arzt etal., 1993; Hebsur etal., 1993; Dymek et al., 1993; Bieler et al., 1993), the nitrides SiN and A1N (Shah et al., 1990; Whittenberger etal., 1990b, 1992a; Moser etal., 1990; Kumar, 1991; Bieler etal., 1992; Arzt et al., 1993; Hebsur et al., 1993), SiC, B4C and other carbides (Whittenberger et al., 1990a; Moser et al., 1990; Nardone etal., 1990; Shah etal., 1990; Chou and Nieh, 1991; Dimiduk etal., 1991; Dunmead etal., 1991; Kumar, 1991), the boride TiB2 (Rigney etal., 1989; Sagib etal., 1990; Moser etal., 1990; Whittenberger etal., 1990b; Alman etal., 1991; Cheng and Cantor, 1992; Kumar etal., 1992a, b; Pregger etal. 1992; Viswanadham et al., 1988; Korinko et al., 1992), and the beryllide TiBe12 (Carbone et al., 1988). At service temperature chemical reactions may occur between the constituent phases of a composite and thus chemical compatibility of the phases to be combined is of primary importance. Microstructural long-term stability is guaranteed only if the constituent phases form stable equilibria with each other. Even then the phase distributions coarsen slowly by ageing processes, i.e. Ostwald ripening occurs which is slower the smaller the interface energies, the mutual solubilities and the diffusivities are (Pitsch and Sauthoff, 1992). Studies of the compatibility of the various strengthening phases with NiAl have shown that A12O3 and TiB2 are sufficiently stable in NiAl (Sagib et al., 1990; Moser et al., 1990; Shah etal., 1990; Chou and Nieh, 1991; Trumble and Ruhle, 1990; Wang and Arsenault, 1991; Korinko and Duquette, 1994). Apart from chemical compatibility, dispersoids may affect the environmental be-

11.4 Nickel Aluminides and Related Phases

havior, e.g. the oxidation resistance, as has been shown recently with respect to an NiAl/AIN composite (Lowell et al, 1990). Besides powder metallurgy methods, directional solidification of eutectic NiAl alloys is used for producing in situ composites. Examples are NiAl-Cr (Walter and Cline, 1970; Kowalski and Frommeyer, 1992). NiAl-Mo (Sauthoff, 1990 a; Subramanian et al., 1990a), NiAl-W (Sauthoff, 1990 a; Saigal and Kupperman, 1991), and NiAl-Re (Mason et al., 1990). In these composites the second continuous phase is disordered, which is expected to improve the ductility of NiAl. However, the obtained ductilities of these composites are still low. Other intermetallic phases are also used for reinforcing an NiAl matrix composite, and an example is given by NiAl containing the Heusler-type phase Ni2AlTi (Whittenberger etal., 1989; Kumar and Whittenberger, 1992) (see also Sec. 11.4.3.6.). A special type of composite is being developed (Nardone etal., 1990; Nardone and Strife, 1991; Nardone, 1992) which combines a continuous ductile phase and a particulate hard and brittle phase with NiAl. In such alloys cracking occurs in the NiAl matrix during deformation at about 0.2% strain, whereas the final fracture strain of the composite is 30-35%. The various NiAl matrix composites usually exhibit an increased strength and brittleness compared to single-phase NiAl. However, there are exceptions to this rule, i.e. a decreased high-temperature strength has been observed for an NiAl-Ni 2 AlTi composite and furthermore, increased toughness has been obtained for an NiAlA12O3 composite (Kumar, 1991). Such toughening by dispersed particles is a wellestablished process for improving the toughness of metals and ceramics (Wiederhorn, 1984; Sigl et al., 1988; Ashby et al.,

723

1989). Another toughening mechanism in ceramics relies on the transformation of dispersed zirconia particles which transform during deformation (Wiederhorn, 1984; Evans and Cannon, 1986; Riihle and Evans, 1989). This transformation toughening with zirconia has been used not only for ceramics, but also for metals and in particular for the intermetallic NiAl, i.e. the fracture toughness of sintered NiAl has been increased from about 14 MN/m 3/2 to 22 MN/m 3/2 by the addition of 20 vol.% ZrO 2 which contained 2 mol% Y 2 O 3 (Barinov et al., 1992; Barinov and Evdokimov, 1993). Present efforts of various NiAl developments are directed at improving the processing for optimizing the mechanical behavior. The physical understanding of the deformation controlling mechanisms is only limited and necessitates much further work. 11.4.4 Other B2 Phases Intermetallic phases with a B2 structure form one of the largest groups of intermetallics, and there is a broad variation of chemical, physical and mechanical properties within this group (Dwight, 1967). Other B2 aluminides are FeAl and CoAl. They are closely related to NiAl since Ni, Fe, and Co can substitute for each other completely in the B2 structure - this was made use of for improving the high temperature strength and creep resistance (Jung et al., 1987; Sauthoff, 1991 a) - and they are also considered for high temperature applications (Vedula and Stephens, 1987 a; Liu etal., 1990; Kumar and Whittenberger, 1992). Besides these aluminides, the B2 structure is adopted by compounds of group VIII metals and group IVA metals, e.g. NiTi, CoTi, and FeCo. These B2 phases, which are oft interest with respect to applications, are briefly overviewed in

724

11 Intermetallics

the following sections with the exception of FeAl which is the subject of Sec. 11.5.3. The classical B2 phase CuZn (P-brass) is addressed in Sec. 11.6. Intermetallics with precious metals have recently been studied with respect to strength, ductility and oxidation resistance (Fleischer, 1992 b; Fleischer and McKee, 1993). Outstanding properties have been found for the B2 phases AlRu and RuSc with melting temperatures of 2060° C and 2200 °C, respectively. AlRu shows very high strength with high work hardening and extensive compressive ductility at room temperature. The effects of constitutional defects and ternary solutes on the hardening behavior of AlRu have been studied in detail (Fleischer, 1993 a, b, d, f)The fracture behavior of AlRu has been studied theoretically by a molecular dynamics simulation (Becquart etal., 1993). The oxidation resistance of AlRu allows the use of uncoated AlRu up to about 1250 °C. Of course the high costs of such intermetallics will not allow the application of these phases as structural materials in the near future (Fleischer, 1992 b). Pdln is noteworthy because of its gold color which makes it promising for dental prosthetics and jewellery (Baker and George, 1992). Pdln has been studied with respect to deformation behavior (Munroe et al., 1991) and diffusion (Koiwa, 1992; Wever, 1992). 11.4.4.1 CoAl The B2 aluminide CoAl is very similar to NiAl with respect to density, thermal expansion behavior, melting temperature and phase diagram (Westbrook, 1956; Whittenberger, 1985; Stephens, 1985; Massalski etal., 1990; Harmouche and Wolfenden, 1987). The thermodynamic properties and point defects have been

studied experimentally and theoretically (Bakker and Ommen, 1978; Chen and Dodd, 1986; Koch and Koenig, 1986; chapter by Inden and Pitsch in Volume 5 of MST). Young's modulus, which is higher than that of NiAl, has been determined as a function of temperature for polycrystalline CoAl with the stoichiometric composition and with various deviations from stoichiometry (Harmouche and Wolfenden, 1985,1986). Diffusion in CoAl is slower with a higher activation energy than in NiAl, and deviations from stoichiometry enhance diffusion as in the case of NiAl (Hagel, 1967). The hardness of CoAl is higher than that of NiAl with a minimum for the stoichiometric composition at temperatures below 800 °C (Westbrook, 1956), i.e. deviations from stoichiometry produce constitutional defects with restricted mobility at lower temperatures which strengthen the CoAl, as in the case of NiAl. The characteristics of slip correspond to those of NiAl (Baker and Munroe, 1990). The hardening of CoAl by constitutional defects and ternary solutes, i.e. Mn, Re, and Ti, has been studied in detail (Fleischer, 1993 c, d, e). At high temperatures above 1000 °C the flow stress and creep resistance are again higher than those of NiAl and reach a maximum at the stoichiometric composition (Hocking et al., 1971; Whittenberger, 1985; Yaney and Nix, 1988) because the constitutional defects, which are produced by deviations from stoichiometry, are mobile at high temperatures and enhance diffusion and thereby creep, as in the case of NiAl. The enhanced creep resistance of CoAl is attributed to lattice friction effects limiting the dislocation mobility in CoAl (Yaney and Nix, 1988). The creep resistance of CoAl is increased significantly by partial substitution of Co by Ni, and still more by Fe (Jung etal., 1987; Whitten-

11.4 Nickel Aluminides and Related Phases

berger, 1987; Sauthoff, 1991a). Further strengthening is achieved by reinforcing CoAl with particulate TiB 2 (Mannan etal., 1990). The high strength of CoAl is correlated with low ductility and toughness, and indeed the fracture toughness is lower than the already low toughness of NiAl (K.-M. Chang etal., 1992). The fracture characteristics, however, are quite analogous to those of NiAl, i.e. the preferred cleavage plane is {110} and the fracture toughness is lowest for directions in the cleavage plane. The differences in strength and brittleness between CoAl and NiAl cannot yet be correlated with differences in the strength and the character of bonding, which have been studied by quantum-mechanical ab initio calculations (Schultz and Davenport, 1992). Because of its high brittleness and because Co is a strategic material, CoAl has not been selected for materials development in the past in spite of its attractive high-temperature strength (Vedula and Stephens, 1987 a). Recent work on CoAl-TiB 2 composites has led to the conclusion that the more damage-tolerant NiAl matrix composites are a better choice for creep applications (Mannan et al., 1990). 11.4.4.2 NiTi The B2 phase NiTi has been used for 30 years as a shape memory alloy for couplings, fasteners, connectors, and actuators in automotive and aerospace industries, electronics, mechanical engineering and medical applications (Schmidt-Mende and Block, 1989; Stockel, 1989; Thier, 1989; Hodgson, 1990; Stoeckel, 1990). NiTi melts congruently at 1310 °C and has an extended homogeneity range (Massalski etal., 1990).

725

NiTi transforms martensitically from the parent B2 phase to a monoclinic martensite phase with an intermediate orthorhombic R phase from which a reversible shape memory effect - with the best shape memory behavior of all shape memory alloys - and pseudoelasticity results (Hwang et al., 1983; Hwang and Wayman, 1983; Nishida and Honma, 1984; Van Humbeek and Delaey, 1989). The characteristics of such shape memory alloys have been discussed in detail with respect to crystallography, microstructure, thermodynamics, kinetics and macroscopic mechanical behavior, and applications have been described (Delaey etal., 1974; Krishnan etal., 1974; Warlimont etal., 1974; Ahlers, 1986; Van Humbeek and Delaey, 1989; Hornbogen, 1991). The crystallography of the martensitic transformation has been analyzed in detail (Shimizu and Tadaki, 1992). The shape memory behavior of NiTi is improved by prior thermomechanical treatments which produce advantageous microstructures and textures (Van Humbeek and Delaey, 1989; Lin and Wu, 1992; Mulder etal., 1993). The transformation temperature is about 110°C for stoichiometric NiTi and is decreased by excess Ni (Hodgson, 1990). Likewise, the transformation temperature is lowered by alloying with Fe and Cr whereas Cu decreases the transformation temperature hysteresis. A widened transformation temperature hysteresis was found recently for NiTi-Nb alloys, which is attractive with respect to coupling and sealing applications (Piao etal., 1992a, b). The tranformation temperature is increased up to 500 °C by the partial substitution of Ni with Pd which, however, affects the ductility (Yang and Mikkola, 1993). The elastic and plastic deformation behavior, including mechanical twinning, of

726

11 intermetallics

NiTi alloys has been studied intensively (see, e.g. Saburi etal., 1984; Matsumoto and Ishiguro, 1989; Goo etal., 1985). Noteworthy is the high ductility of NiTi which allows deformations up to 70% (Van Humbeek and Delaey, 1989). Furthermore, the yield strength of NiTi depends on the temperature in an anomalous way, i.e. it increases with rising temperature to reach a maximum, and only above this peak temperature does normal softening occur (Takasugi etal., 1991a). This strength anomaly is also shown by other, closely related B2 phases, i.e. CoTi, CoZr, and CoHf (see next section), but not by FeTi and not by the aluminides NiAl, CoAl and FeAl, i.e. the transition from normal behavior to anomalous behavior is obviously related to a critical number of valence electrons per atom (Takasugi etal., 1991 e). The general characteristics of the mechanical behavior of such shape memory alloys and the mutual interaction of deformation and phase transformation have been discussed in detail, and the thermodynamics and kinetics of the involved processes have been addressed (Delaey etal., 1974; Krishnan etal., 1974; Warlimont et al., 1974; Hornbogen 1991). Processing of NiTi alloys is difficult and has to be done with care to obtain advantageous properties (Van Humbeek and Delaey, 1989; Hodgson, 1990). Melting must be done in a vacuum or in inert atmospheres because of the reactivity of Ti, and thus vacuum-induction melting, plasmaarc melting, and electron-beam melting are all used commercially. Hot working is possible without problems. Processing at low temperatures is difficult because of extremely high work hardening, and thus intermediate annealing treatments are required for cold rolling, machining, etc. The ductility problem of the (Ni,Pd)Ti alloys with a high transition temperature is re-

lieved by alloying with B (Yang and Mikkola, 1993). Joining by welding, brazing or soldering is generally difficult. The excellent shape memory behavior of the NiTi alloys is accompanied by excellent corrosion resistance - comparable to stainless steels - which make these alloys the only shape memory alloys suitable for implantation in human bodies (Van Humbeek and Delaey, 1989). Recently NiTi has been used as a strengthening phase in Al matrix composites which experience an additional strengthening effect from the shape memory effect in NiTi (Furuya etal., 1993; Yamada etal., 1993). 11.4.4.3 FeTi, CoTi, CoZr, and CoHf FeTi with a B2 structure is one of the most promising intermetallic phases for use as a hydrogen storage material and related applications (Reilly, 1979; Nathrath, 1986; Lynch, 1991). Dissolved hydrogen orders to form a substructure in which the B2 structure becomes rhombic, i.e. a stable hydride is formed in a reversible way (Somenkov and Shilstein, 1979; Reilly, 1979). FeTi with excess Ti is more easily hydrogenized then stoichiometric FeTi since Ti activates the absorption and desorption of hydrogen catalytically (Sicking and Jungblut, 1983; Amano et al., 1984). Second phases, including oxides, shorten the incubation time for hydrogenization (Nagai et al., 1987; Amano etal., 1983), whereas oxygen may poison the FeTi surface at room temperature (Schlapbach etal., 1979). The electronic structure has been studied theoretically (Bose et al., 1991). CoTi is also of interest for use as a hydrogen storage material (Izumi, 1989), as well as the ternaries (Fe,Co)Ti (Boulghallat and Gerard, 1991) and (Fe,Ni)Ti (Bershadsky et al., 1991) which show more advantageous storage properties than the binaries.

11.4 Nickel Aluminides and Related Phases

The plastic deformation behavior of CoTi has been studied in detail since the B2 phases CoTi, CoZr, and CoHf exhibit an anomalous temperature dependence of the yield strength, which is also shown by NiTi, but not by FeTi and the aluminides NiAl, CoAl, and FeAl (Nakamura and Sakka, 1988; Takasugi and Izumi, 1988; Takasugi et al., 1990c, 1991 a, e; Yoo et al., 1990; Yoshida and Takasugi, 1991b). Slip occurs by <100> dislocations at low and intermediate temperatures as in the case of NiAl, whereas at higher temperatures above the peak strength temperature the activation of <111> dislocations has also been observed (Takasugi etal., 1992b). The elastic moduli have been measured (Yasuda et al., 1991 b). The transition from the anomalous temperature dependence of the yield strength in NiTi, CoTi, CoZr, and ZrHf to the normal one in FeTi and the aluminides NiAl, CoAl, and FeAl is related to the change in the number of valence electrons per atom, i.e. to changes in the atomic bonding, but a clear and complete understanding has not yet been reached (Yoo etal., 1990; Takasugi etal., 1991 e, 1992 b). It is noted that alloying CoTi with Co Si results in the formation of the Heusler-type compound Co2TiSi with an L2X structure, which is strongly ferromagnetic with a Curie temperature far in excess of room temperature though both CoTi and CoSi are nonmagnetic (see also Sec. 11.4.5). 11.4.4.4 FeCo FeCo - in particular with a 2 % V addition - has found widespread use as a magnetically soft material with a high saturation induction, e.g. as electromagnet pole tips, and the respective alloys are known as Permendur alloys (Chen, 1961; Kouvel, 1967; Dietrich, 1990). These alloys also exhibit a high, positive magnetostrictive co-

727

efficient, which has made them useful in sonar transducers and in extremely accurate positioning devices. The B2 phase FeCo is stable only at temperatures below 730 °C and has a broad range of homogeneity (Massalski et al., 1990). At 730 °C there is an order-disorder transition with a 0.2 % change in volume from B2 to A2 (b.c.c), and a further structure change from A2 to Al (f.c.c.) occurs at about 985 °C (Buckley, 1975). The ordering reaction proceeds heterogeneously or homogeneously depending on the temperature (Cahn, 1993). The ordering process has been studied in detail and the thermodynamics of the binary Fe-Co phase diagram, as well as the ternary Fe-Co-Al phase diagram, have been analyzed by model calculations (Rajkovic and Buckley, 1981; chapter by Inden and Pitsch in Volume 5 of MST). The ordering kinetics are affected by alloying additions and in particular alloying with 2 at.% V reduces the ordering rate significantly, which allows freezing of the high-temperature, disordered structure by quenching to low temperatures. It is noted that the addition of more than 2 % V leads to alloy decomposition, as is indicated by the Fe-Co-V phase diagram (Mahajan etal., 1974; Raynor and Rivlin, 1983). A still larger effect on the ordering kinetics is produced by alloying with Cr (Alekseyev et al., 1977). Below the critical temperature of ordering, the degree of long-range order in equilibrium increases continuously from 0 at 730 °C to 0.9 at 500 °C to reach 1 - corresponding to perfect order - asymptotically at low temperatures (Stoloff and Davies, 1964). The mechanical behavior of FeCo has been studied intensively since the orderdisorder transition allows a direct study of the effects of ordering on the mechanical properties including creep, fatigue and

728

11 Intermetallics

fracture (Stoloff and Davies, 1964; Boettner et al., 1966; Jordan and Stoloff, 1969; Marcinkowski, 1974 b; Pitt and Rawlings, 1983; Kawahara, 1983 a, c; Delobelle and Oytana, 1983; Miiller, 1986; Stoloff et al, 1992; Zhao et al., 1993). At room temperature FeCo is ductile - with 15 % uniform elongation and subsequent necking - in the disordered state, i.e. after quenching from above the ordering temperature, whereas in the ordered state, i.e. after annealing below the ordering temperature, it shows only a little ductility with 5% uniform elongation and no necking (Stoloff and Davies, 1964). Thus cold rolling is difficult in the ordered state (see Miiller, 1986). Furthermore, FeCo is subject to hydrogen embrittlement at room temperature (Liu and Stoloff, 1993). Plastic deformation at room temperature occurs by the movement of <111> superlattice dislocations which are pairs of superpartials (Baker and Munroe, 1990; Yamaguchi and Umakoshi, 1990). The flow stress increases with decreasing degree of order, i.e. with rising temperature, to reach a maximum at about 700 °C, just below the ordering temperature where FeCo is only partially ordered (Stoloff and Davies, 1964). This behavior results from a decoupling of the superpartials of the superlattice dislocations since the coupling force decreases with increasing disorder. Above the flow stress peak temperature normal softening is observed which is due to thermal activation. In other words, FeCo, like other intermetallics, exhibits an anomalous positive temperature dependence of the flow stress, which, however, is not related to special dislocation mechanisms as in the case of Ni 3 Al, but is related to incomplete atomic order, i.e. to structure changes. Ordered FeCo-2V exhibits a more homogeneous deformation and a higher

strain hardening rate than disordered FeCo-2V at all temperatures (Jordan and Stoloff, 1969). The grain size dependence of the yield stress and the fracture stress is described by the usual Hall-Petch relationship. The fatigue resistance is improved by ordering with only a small effect of the environment (Boettner et al., 1966; Stoloff et al., 1992). The low ductility can be increased by adding Ni as a fourth element, which changes the microstructure, and still more by thermomechanical treatments which lead to a Hall-Petch-type dependence of ductility on grain size or subgrain size (Pinnel etal. 1976; Pitt and Rawlings, 1983). Even cold rolling can reduce the brittleness (Kawahara, 1983 c). The brittle-to-ductile transition temperature in the ordered state can be reduced significantly by deviations from stoichiometry which decrease the critical temperature of ordering and introduce configurational disorder (Konoplev and Sarrak, 1982). Positive effects on cold workability have been produced - apart from alloying with V or Cr which both slow down the ordering kinetics - by alloying with Nb, Ta, Mo, W, and C whereas Be, Ti, Zr, Mn, Cu, Ag, Au, B, Al, and Si are ineffective (Kawahara, 1983 a, b). The physical understanding of such effects is still to be clarified. Finally, it may be noted that alloying Fe and Co with precious metals leads to the phases FePt, CoPt, and FePd with an L l 0 structure, which is an ordered f.c.c. structure with tetragonal distortion (see Fig. 11-1). These phases have been of importance for applications as permanent magnetic materials, and in particular CoPt, which is used in the only partially ordered state, excels because of isotropy, ductility, easy machinability and resistance to corrosion and high temperatures (Jellinghaus, 1936; Kouvel, 1967; Chin and Wernick,

11.4 Nickel Aluminides and Related Phases

1986; Fiepke, 1990; Leroux et al., 1991; Watanabe, 1991). However, CoPt alloy magnets are seldom used now since these alloys have been replaced by the rare earth magnets with superior magnetic properties. 11.4.5 Heusler-Type Phases The classic Heusler-type alloys are Cu2AlMn and Cu 2 SnMn (Dwight, 1967; Kouvel, 1967) which will be addressed in Sec. 11.6. Their unique feature is their ferromagnetism, since the constituent elements are not ferromagnetic. These phases crystallize with the L2A structure which is derived from the B2 structure (see Fig. 11-1), and which is also adopted by a series of other ternary phases (Dwight, 1967). Outstanding examples are Ni 2 AlX and Co2AlX phases with X = Ti, Zr, Hf, Nb, or Ta which are of interest for high-temperature applications, and some of them have been studied with respect to hightemperature deformation (Strutt et al., 1976; Umakoshi et al., 1985, 1986; Sauthoff, 1989; Takeyama et al., 1991). In particular, Ni2AlTi has been used for reinforcing NiAl (see Sec. 11.4.3.6 under the heading Structural Alloys) and a twophase Ni 2 AlTi-NiAl showed a creep resistance which was higher than that of singlephase NiAl and even Ni2AlTi, and reached that of a high-strength Ni-base superalloy (Polvani et al., 1976). The high strength is correlated with brittleness which has handicapped respective materials developments until now. Only recently Ni 2 AlTaNiAl has been identified as a more promising system because of the improved ductility (Pak et al., 1988). 11.4.6 Nickel-Molybdenum Phases Various intermetallic phases are formed in the binary Ni-Mo system (Brooks et al.,

729

1984). The most Ni-rich phase Ni 4 Mo has the body-centered tetragonal crystal structure D l a and is stable only up to about 880 °C with peritectic decomposition. The ordered atom distribution can be established in various ways in this structure which gives rise to a domain structure with various types of domain interfaces. Ni 3 Mo is orthorhombic and is stable up to about 910 °C. This phase, however, can also be formed with the metastable, tetragonal D0 22 structure, which is common to the trialuminide Al3Ti and related phases (see Sec. 11.3.3). NiMo is another orthorhombic phase and decomposes peritectically at 1362°C. Finally, there is the metastable, orthorhombic Ni 2 Mo phase which is formed as an intermediate phase during ordering. These phases are of general interest with respect to alloying Ni-base alloys since Mo is an important alloying element for superalloys. Corresponding multinary phase diagrams have been studied experimentally and theoretically (see, e.g., Brooks et al., 1984; Chakravorty and West, 1986; Kodentzov et al., 1988; Enomoto et al., 1991), in particular with respect to equilibria with the aluminide Ni 3 Al which forms the basis of the advanced aluminides (see Sec. 11.4.1). Special interest is directed at Ni 4 Mo since Ni-Mo alloys with about this composition (20 at.% Mo) and additions of further elements - in particular Fe, Cr, and C - have been of industrial significance since the early 1930s because of their excellent corrosion resistance in nonoxidizing environments and their high strength at high temperatures, as is exemplified by the Hastelloy B alloys (Brooks et al., 1984). In these alloys ordered Ni 4 Mo may form on slow cooling from temperatures above 900 °C, where the f.c.c. solid solution is stable, or on ageing at 500-850 °C.

730

11 Intermetallics

Ni 4 Mo formation in Ni-Mo alloys leads to both hardening and embrittlement (in tension) with fracture occurring along grain boundaries. In compression such embrittled alloys can still be deformed plastically and furthermore, monocrystalline ordered Ni 4 Mo shows significant ductility even in tension (Brooks et al., 1984; Kao et al., 1989). The physical properties of ordered Ni 4 Mo have been determined and the deformation behavior has been studied in detail (Brooks etal., 1984). The mechanical properties are a sensitive function of the microstructure. The microstructure development is controlled by the kinetics of the decomposition and ordering reactions during cooling and ageing of these Ni-Mo alloys (see, e.g. Brooks and Cao, 1992; Cahn, 1992). The low ductility of hypostoichiometric ordered Ni 4 Mo can be improved significantly by microalloying with B, whereas there is only a small effect of B on the stoichiometric phase (Tawancy, 1991).

11.5 Iron Aluminides and Related Phases An overview of the Fe-Al phases is given by the binary Fe-Al phase diagram (Massalski et al., 1990). The Al-rich phases with melting temperatures of about 1150 °C are of interest for Al alloys and are not addressed here. The remaining phases Fe3Al and FeAl, and related phases, have been selected for materials developments because of their outstanding physical, mechanical, and chemical properties and are the subjects of the following sections. 11.5.1 Fe3Al Fe3Al is formed on cooling by ordering reactions in the solid state that transform the b.c.c. disordered solid solution, which

is stable above about 800 °C, first into the FeAl phase with a B2 structure, which is stable between about 800 °C and 550 °C and which is the subject of Sec. 11.5.3, and then into Fe3Al with the D0 3 structure (see Fig. 11-1). The critical temperature of ordering is much lower than the melting temperature (Massalski et al., 1990) which indicates less strong bonding between unlike atoms compared with, e.g. the nickel aluminides. The ordering process and its kinetics have been studied in detail both experimentally and theoretically for the determination of the field of existence and the equilibria of Fe3Al in the Fe-Al phase diagram (see, e.g. Koster and Godecke, 1980; Godecke and Koster, 1985; Wachtel and Bahle, 1987; Inden and Pepperhoff, 1990; Vennegues etal., 1990; Maziasz et al., 1993; chapter by Inden and Pitsch in Volume 5 of MST). Recently a transient B32 phase has been observed during ordering (Gao and Fultz, 1993). On the Fe-rich side there is a two-phase equilibrium between Fe3Al and b.c.c. Fe-Al solid solution which shows short-range order. On the Al-rich side there is a second order transition from Fe3Al with a D0 3 structure to FeAl with a B2 structure without a twophase equilibrium. The generally accepted phase diagram with the D03/B2 transition has been questioned recently on the basis of neutron scattering results (Hilfrich etal., 1991) which, however, are not unambiguous (Inden, 1993). The critical temperature of ordering, which is about 550 °C in the binary case (Massalski et al., 1990), can be shifted to higher temperatures by as much as 250 °C by alloying with a third element, in particular Cr, Mo, Mn, Ti, or Si (Mendiratta and Lipsitt, 1985; Fortnum and Mikkola, 1987; Longworth and Mikkola, 1987; Prakash et al., 1993). Thermal vacancies are formed readily because of a low formation enthalpy,

11.5 Iron Aluminides and Related Phases

which is even lower than the migration enthalpy, and consequently there is a high thermal-equilibrium vacancy concentration in contrast to e.g. Ni 3 Al and pure metals (Schaefer etal., 1990, 1992). The resulting activation energy of diffusion, which has been determined by studying the ordering kinetics, varies strongly with the Al content (Vennegues et al., 1990). The diffusion coefficients for B, Ni, V, and Ti in Fe3Al with a D0 3 or a B2 structure have been determined recently (Hasaka et al., 1993 a, b). Fe3Al exhibits an outstanding, high magnetic permeability which makes it useful as a magnetic material (Nachman and Buehler, 1954). The magnetic behavior can still be improved by partial substitution of Al by Si, which does not change the crystal structure. This gave rise to the development of the Sendust alloys which are based on Fe3(Al,Si) and which have found widespread use as magnetic head materials in recorders (Yamamoto and Utsushikawa, 1977; Yamamoto, 1980; Watanabe etal., 1984; Brock, 1986). Besides this application as a functional material, Fe3Al is now also being considered for structural applications since it not only shows a higher strength than comparable iron alloys, but also a high corrosion resistance in oxidizing and sulfidizing environments (McKamey etal., 1991). Problems for application are the low ductility at ambient temperatures and the strength drop above the ordering temperature. The plastic deformation, including twinning, of Fe3Al and the effect of atomic order on strength and ductility has been studied in detail (Marcinkowski and Brown, 1961; Stoloff and Davies, 1964, 1966; Morgand etal., 1968; Leamy and Kayser, 1969; Marcinkowski, 1974 a; Hanada et al., 1981 a, b; Park et al., 1991; Ehlers and Mendiratta, 1984; Schroer et al.,

731

1991; Shindo etal., 1991). Ductility depends sensitively on prior processing (McKamey and Pierce, 1993), and plastic deformation may reduce the degree of longrange order (Dadras and Morris, 1993). A special feature is the anomalous temperature dependence of the flow stress, which reaches a maximum slightly below the ordering temperature because of partial disorder. This behavior is thought to be due to the decrease in order, as in the case of FeCo which has been discussed in Sec. 11.4.4.4. However, new results have shown that the yield stress anomaly is a result of the complex temperature dependence of the dislocation mobilities (Schroer et al., 1993). Dislocation line energies and line tensions as well as elastic moduli have been determined for both Fe3Al and Fe3(Al,Si) (Kotter et al., 1989). Plastic deformation of Fe3Al under special experimental conditions gives rise to pseudoelasticity and a shape memory effect (Kubin et al., 1982; Nosova et al., 1986). Fe3Al has been alloyed successfully with Ni and Cr to improve the ductility (Horton etal., 1984; Morris etal., 1993a). Good combinations of strength and ductility have been obtained by thermomechanical treatments which resulted in large amounts of B2 phase (Sun et al., 1993). If the alloying additions surpass the solubility limits, precipitation of second phases, e.g. a Heusler-type phase with a cubic L2X structure or a Laves phase with a hexagonal C14 structure, occur which may be used for improving the strength of Fe3Al (Dimiduk etal., 1988; Ranganath etal., 1991; Maziasz and McKamey, 1992). The respective Fe-Al-Ni alloys are closely related to the NiAl-based Ni-Fe-Al alloys which have been addressed in Sec. 11.4.3.6 (under the heading Structural Alloys) and which have given rise to corresponding materials developments.

732

11 Intermetallics

In view of structural applications, Fe3Al has been studied with respect to creep (Lawley e t a l , 1960; Davies, 1963 a; Prakash et al., 1991 b), fatigue (Fuchs and Stoloff, 1988; Prakash etal., 1991b; Castagna and Stoloff, 1992; Stoloff etal., 1993), fracture (Sainfort et al., 1963; Horton etal., 1984; Mendiratta etal., 1987a; Shan and Lin, 1992), and oxidation resistance, including Fe3Al-based alloys and composites (Cathcart, 1985; DeVan, 1989; Tortorelli and DeVan, 1992; Nourbakhsh etal., 1993). The ductility of Fe3Al is a very sensitive function of the environment, i.e. Fe3Al is subject to environmental embrittlement which is due to the presence of moisture in the surrounding atmosphere (McKamey and Liu, 1990; Kasul and Heldt, 1991; McKamey etal., 1991; Sanders etal., 1991; Liu, 1992; Hippsley and Strangwood, 1992; Shan and Lin, 1992; Vyas et al., 1992; Lin et al., 1992; Castagna etal., 1993; Stoloff and Duquette, 1993). Even very low partial pressures of water vapor are sufficient for embrittlement (McKamey and Lee, 1993). These embrittling effects are absent only for such low Al contents that D0 3 long-range ordering no longer occurs (Sikka et al., 1993 b). In spite of the mentioned shortcomings of Fe 3 Al, this iron aluminide is of great industrial interest, and various alloy developments have been initiated which rely on composition optimization and thermomechanical treatments for improving the mechanical behavior (Liu etal., 1990; Esslinger and Smarsly, 1991; Baker and George, 1992; Liu and Kumar, 1993; Viswanathan et al., 1993; Sikka etal., 1993a). Recently an Fe3Al matrix composite with reinforcing A12O3 dispersoids has been produced with good strength and oxidation resistance, and with sufficient ductility (Suganuma, 1993).

As to processing, both powder metallurgy and ingot metallurgy - including arc melting, vaccum-arc remelting, induction melting in air and vacuum, and electroslag remelting - have been studied (Sanders et al., 1991; Sikka et al., 1991, 1993 a; Sikka, 1991; Gieseke etal., 1993). These developments aim at applications as automotive exhaust systems and resistance heating elements, which are close to commercialization. Other potential applications are steam turbine discs and superheater tubes in power plants, hot gas filters in coal gasification plants, and chemical processing equipment. Recently combustion synthesis of Fe 3 Al, which has been used to join SiC composites, has been proposed for coating structural steels in view of the corrosion and wear resistance of Fe3Al (Wright etal., 1993). 11.5.2 Fe 3 AlC x and Related Phases

The dissolution of carbon in Fe3Al produces the so-called K phase which is usually given as Fe 3 AlC x with x « 0.66 for simplicity (Huetter and Stadelmaier, 1958; Goldschmidt, 1967 a, 1969; Nowotny, 1972 b) though it shows an extended range of homogeneity corresponding to Fe 3 _ y Al 1 + J C x with - 0.2 < y < + 0.2 and 0.42 < x < 0.71 (Andryushchenko et al., 1985; Palm, 1990). Its crystal structure L'l 2 is related to the Ll 2 structure of Ni 3 Al since the Fe and Al atoms occupy the sites of an LI 2 sublattice in this phase, whereas the C atoms occupy the centers of the L l 2 unit cells (see Fig. 11-1). Thus it may be regarded as an intermetallic L l 2 phase which is stabilized by the interstitial C. On the other hand, it is an interstitial compound and is usually regarded as a complex carbide with a crystal structure corresponding to the perovskite-type structure E2i (with x = 1). F e ^ l C , , is hard

11.5 Iron Aluminides and Related Phases

and brittle, it has an appreciable range of homogeneity, tranformations are not known, and - in contrast to Morral (1934), and Huetter and Stadelmaier (1958) - it is not ferromagnetic (Lohberg and Schmidt, 1938; Parker et al., 1988). The Fe-rich corner of the Fe-Al-C phase diagram has been studied recently in order to clarify the phase equilibria of Fe 3 AlC x with the Fe solid solution on the one hand and with the graphite on the other (Palm, 1990). It was found that Fe 3 AlC x melts at about 1400 °C depending on its composition, and its field of existence in the ternary phase diagram is shifted to higher C contents with decreasing temperature which makes the preparation of monophase Fe 3 AlC x difficult. Other recent work has been directed at the mechanical behavior of Fe 3 AlC c (Jung and Sauthoff, 1989 b; Wunnike-Sanders, 1993). It can be shown that the deformation behavior of Fe 3 AlC x corresponds to that of Ni3Al with dissolved C, which is expected because of the structural similarity. The variation of the C content of Fe3AlCJC leads to the precipitation of either oc-Fe or graphite, which both affect the deformation behavior, i.e. flow stress, causing the creep resistance and the brittle-toductile transition temperature to increase in the sequence Fe 3 AlC x + a-Fe, Fe 3 AlC x , Fe3A1CX + graphite. Thus there are some prospects for alloy development by optimizing composition and phase distribution. The fracture toughness increases with increasing deformation rate, which indicates the presence of environmental embrittlement (Vehoff, 1993). Finally, it is noted that such perovskitetype phases are also formed in other alloy systems by alloying with carbon as well as with oxygen, and that such phases are considered for strengthening intermetallic alloys (Goldschmidt, 1967 a; Nowotny,

733

1972b; Kassem and Koch, 1991; Ellner et al., 1992; Nemoto et al., 1992). Recently multiphase Co-Co 3 AlC alloys have been prepared with good combinations of strength and ductility (Hosoda et al., 1993). 11.5.3 FeAl FeAl is closely related to NiAl (Sec. 11.4.3) since both phases show the B2 structure and complete mutual miscibility (Bradley and Taylor, 1938; Bradley, 1951; Baker and Munroe, 1990). However, FeAl melts incongruently in contrast to NiAl, and its melting temperature is lower, which indicates lower stability and atomic bonding strength. The binary Fe-Al phase diagram and the phase equilibrium with Fe3Al have already been addressed in Sec. 11.5.1. It has to be noted that a subdivision of the field of existence for the B2 phase in the binary Fe-Al phase diagram was proposed according to observed changes in the physical properties (Koster and Godecke, 1980; Godecke and Koster, 1985, 1986). However, these changes are not related to changes in crystal symmetry or atomic order, and thus the understanding of this subdivision of the B2 phase field is still unclear. FeAl has been studied in detail with respect to density (Baker and George, 1992), elastic behavior - both experimentally and theoretically by quantum-mechanical, ab initio calculations (Marcinkowski, 1974 a; Harmouche and Wolfenden, 1985, 1986; Godecke and Koster, 1986; Fu and Yoo, 1991, 1992a; Yoo and Fu, 1991), point defects (Paris etal., 1975; Ho and Dodd, 1978; Koch and Koenig, 1986) and diffusion (Hagel, 1967; Cheng and Dayananda, 1979; Akuezue and Whittle, 1983; Bakker, 1984; Vogl and Sepiol, 1994). Most work on FeAl has been directed at the mechani-

734

11 Intermetallics

cal behavior and in particular, at the strength and ductility (Sainfort et al., 1963; Morgand etal., 1968; Marcinkowski, 1974a; Crimp etal., 1987; Baker and Munroe, 1990; Baker etal., 1991; Lefort etal., 1991; Nagpal and Baker, 1991; Prakash etal., 1991b; Tassa etal., 1991; Baker and Nagpal, 1993), and special attention has been paid to the analysis of dislocation behavior (Strutt and Dodd, 1970; Marcinkowski, 1974 a; Umakoshi and Yamaguchi, 1980; Mendiratta et al., 1984; Rudy and Sauthoff, 1986; Mendiratta and Law, 1987; Feng and Sadananda, 1990; Munroe and Baker, 1991; Takahashi and Umakoshi, 1991). In particular, <100> slip, which is characteristic for NiAl, has been reported only at high temperatures, whereas <111> slip has been observed at lower temperatures with a transition temperature that increases with decreasing Al content (Yamagata and Yoshida, 1973; Umakoshi and Yamaguchi, 1980, 1981; Mendiratta et al., 1984). The <100> dislocations are perfect single dislocations, whereas the <111> dislocations are superlattice dislocations consisting of two partial dislocations with an antiphase boundary (APB) in between. The APB energy, which contributes to the superlattice dislocation energy, decreases with decreasing Al content according to quantum-mechanical, ab initio calculations (Freeman et al., 1991). The transition from <111> slip to <100> is correlated with a yield strength peak and a ductility drop at about 450 °C for off-stoichiometric FeAl (Xiao and Baker, 1993; Guo et al., 1993; Yoshimi and Hanada, 1993). It is noted that strength and hardness at low temperatures do not reach a minimum at the stoichiometric composition, which is in contrast to the behavior of the isomorphous phases NiAl and CoAl (see Sec. 11.4.3.3), and which is not yet under-

stood (Westbrook, 1956; Crimp and Vedula, 1991). Recently it has been found that the room temperature hardness increases with increasing vacancy concentration, and the latter increases with increasing Al content through the stoichiometric composition (Chang et al., 1993). The vacancy formation energy is low with a strong tendency for vacancy clustering according to theoretical ab initio calculations (Fu et al., 1993 a). Single-crystalline FeAl shows some ductility in compression at low temperatures whereas the tensile ductility is not known (Yamagata and Yoshida, 1973; Baker and Munroe, 1990). The tensile ductility at low temperatures, as well as the compressive ductility, of polycrystalline FeAl is practically nil at the stoichiometric composition - which again is in contrast to NiAl - and it increases with decreasing Al content, i.e. increasing deviation from stoichiometry (Baker and Gaydosh, 1987; Mendiratta etal., 1987b; Schmidt etal., 1989). It has to be emphasized that the reported ductilities of off-stoichiometric FeAl are obtained only with well-annealed specimens, i.e. if the low-temperature deformation is preceded by anneals with very low cooling rates (Schmidt etal., 1989; Nagpal and Baker, 1990 a). This behavior is supposed to be due to retained excess vacancies which increase the hardness and yield strength and decrease the ductility, and which are not equilibrated after high-temperature anneals with normal cooling rates (Schmidt etal., 1989; Nagpal and Baker, 1990 a; Crimp and Vedula, 1991; Kong and Munroe, 1993; Baker and Nagpal, 1993). Vacancy equilibration seems to be a very slow process and thus anneals of several hundred hours at about 400 °C are recommended for producing ductility in off-stoichiometric FeAl (Nagpal and Baker, 1990 a).

11.5 Iron Aluminides and Related Phases

In view of the similarities and differences between FeAl and NiAl, the question again arises in what way the strength and character of atomic bonding changes with composition, and in what way such changes are correlated with the observed changes in the mechanical behavior. However, in spite of the much improved knowledge and understanding of the electronic distribution in these B2 aluminides simple correlations between strength and character of bonding and mechanical behavior are not visible (Schultz and Davenport, 1992), which means that there may be no such simple correlations. FeAl has been alloyed with various elements to improve the mechanical behavior (see e.g., Munroe and Baker, 1990; Lefort et al., 1991; Prakash et al., 1991 a, b; D. Li et al., 1993). Alloying with Ni has initiated promising alloy developments which have been discussed in Sec. 11.4.3.6. It is noted that the above-mentioned transition from <111> slip to <100> slip at elevated temperatures is also effected by alloying with Ni (Patrick et al., 1991). Microalloying with B increases the strength and toughness of FeAl significantly, whereas the brittle-to-ductile transition temperature is reduced only slightly (Crimp and Vedula, 1986; Webb et al., 1993 b; Schneibel et al., 1993 b; Baker, 1993). FeAl is subject to environmental embrittlement, i.e. its ductility depends sensitively on the test environment (Liu and George, 1991; Liu et al., 1991; Lynch et al., 1991; Shea et al., 1991; Klein and Baker, 1992; Schneibel and Jenkins, 1993; Nagpal and Baker, 1991; Klein et al., 1993). This embrittling effect is due to the dissolution of hydrogen in FeAl at the crack tip and is reduced by increasing the deformation rate sufficiently (Nagpal and Baker, 1991; Liu and George, 1991; Liu et al., 1991; Lynch etal., 1991; Shea et al., 1991; Klein and

735

Baker, 1992; Schneibel and Jenkins, 1993; Schneibel etal., 1993b). Dispersoids have been used successfully for strengthening FeAl and respective developments of particulate composites have been initiated (see, e.g. Mendiratta etal., 1987b; Moser et al., 1990; Decamps et al., 1991; Kumar, 1991; Prakash etal., 1991b; Baker and George, 1992; Schneibel et al., 1992b; Xu etal., 1993). Apart from strength and ductility, FeAl has been studied with respect to creep, which has been discussed together with NiAl in Sec. 11.4.3.4, fracture - both experimentally and theoretically by quantum-mechanical, ab initio calculations (Baker and Gaydosh, 1987; Baker et al., 1991; K. M. Chang etal., 1992; Prakash etal., 1992; Yoo and Fu, 1992; Schneibel and Jenkins, 1993), fatigue (Prakash et al., 1991 b; Stoloff, 1992), and corrosion which is oxidation in most cases (Cathcart, 1985; Schmidt etal., 1989; Smialek etal., 1989, 1990 a; Nesbitt and Lowell, 1993). In view of the reported advantageous properties and the comparatively low density, FeAl is regarded as potential structural material for high temperature applications and respective processing developments are making use of both ingot metallurgy and powder metallurgy (Stephens, 1985; Vedula etal., 1985; Vedula and Stephens, 1987a, b; Vedula, 1989, 1991; Liu and Kumar, 1993). The published results are quite promising and FeAl alloy compositions have been identified, e.g. for good weldability (Maziasz et al., 1992) and for high corrosion and wear resistance in aggressive gasification environments (Magnee et al., 1991).

736

11 Intermetallics

11.6 Cu-Base Phases

11.6.1 CuZn

Various intermetallic Cu-base phases have been and are major constituents of many Cu-base alloys - in particular bronzes, i.e. Cu-Sn alloys, and brasses, i.e. Cu-Zn alloys - which have been in use since the beginning of metallurgy (GmelinInstitut, 1955; Westbrook, 1977; Flinn, 1986), and examples are given in Table 111. Various Cu-Sn and Cu-Zn phases have been reviewed with respect to their basic properties (Schubert, 1967; Dwight, 1967; Guillet and Le Roux, 1967; Hagel, 1967). Alloying with a third element results in a still larger multitude of phases. An outstanding example is given by the Heuslertype phases Cu 2 MnAl and Cu 2 MnSn with an L2 t structure (see Fig. 11-1), which are ferromagnetic though the constituent metals are not ferromagnetic (Dwight, 1967; Kouvel, 1967). It is noted that Cu 2 MnAlbase alloys were used for applications as fruit knives at the beginning of this century (Heusler, 1989) - see Table 11-1. Another similar Cu phase is Ni 2 CuSn with an L21 or a D0 3 structure which is considered for structural applications (Kratochvil, 1990; Kratochvil et al, 1992). Some of these binary and ternary phases have been studied to a larger extent because of special physical and/or mechanical properties. Of particular interest are the Cu phases CuZn, Cu^ t + x 3jc Zn^ x _ 2 ^Al^, and (Cu,Ni)3Al, on which the Cu-Zn-Al and Cu-Al-Ni shape memory alloys are based and which are the subject of the following sections. In addition, the Cu-Au phases Cu 3 Au and CuAu and the Cu-Sn phases Cu3Sn and Cu 6 Sn 5 will be addressed, which are important constituents of Cu-Au alloys and amalgams for dental restorative applications.

The major or exclusive constituent of yellow brass is (i brass which is the intermetallic CuZn phase. It exhibits an A2 structure at high temperatures and a B2 structure at low temperatures, i.e. there is an order-disorder transition at about 460 °C (Flinn, 1986; Massalski et al., 1990). Its range of homogeneity - between about 40 and 50 at.% Zn at higher temperatures - depends sensitively on temperature and does not include the stoichiometric 50 at.% composition at intermediate temperatures. This order-disorder transition has been used to study the effect of ordering, e.g. on elastic behavior (Westbrook, 1960 a; Guillet and Le Roux, 1967), diffusion (Girifalco, 1964; Hagel, 1967; Wever et al., 1989; Wever, 1992), recrystallization (Cahn, 1991), and hardness (Westbrook, 1960 a). The deformation behavior of B2 CuZn has been analyzed in detail with respect to the operating slip mechanism (Rachinger and Cottrell, 1956; Fleischer, 1988; Baker and Munroe, 1990), and the dislocation energies and the antiphase boundary energies which control the slip characteristics have been determined (Marcinkowski, 1974a; Beauchamp et al., 1992; Beauchamp and Dirras, 1993). A special feature of the deformation behavior is the anomalous positive temperature dependence of the yield stress which is also known for other intermetallics. It has to be noted that the case of CuZn was the first to be observed and analyzed which led to the understanding of the variation of yield stress with degree of long-range order (Brown, 1959; Saka et al., 1985; Dirras et al., 1992; Matsumoto and Saka, 1993). With increasing temperature the degree of order decreases continuously until the critical temperature of disordering is reached and

11.6 Cu-Base Phases

the second-order transition from B2 order to A2 disorder occurs. Correspondingly, the yield stress reaches a maximum below the critical temperature of disordering. Apart from yielding, creep has been studied in detail (Westbrook, 1960 a; Stoloff and Davies, 1966; Strutt and Dodd, 1970; Delobelle and Oytana, 1983; Hong and Weertman, 1986), as well as hot deformation processing (Padmavardhani and Prasad, 1991), and cyclic deformation (Kawazoe et al, 1989). Besides the order-disorder transition, CuZn in the B2 state can transform martensitically after sufficiently rapid cooling (Ahlers, 1986). Various types of martensite, which differ by the stacking sequence of the close-packed lattice planes, can be formed depending on composition, i.e. electron concentration. It is noted that this martensitic transformation is related to an exceptionally low value of the elastic shear constant which additionally shows an anomalous, positive temperature dependence (Verlinden and Delaey, 1988 a). The martensitic transformation of B2 ordered CuZn can be induced by external elastic stress which gives rise to a shape memory effect for not too small grain sizes (Schroeder et al., 1976). The characteristics of such shape memory alloys have been discussed in detail with respect to crystallography, microstructure, thermodynamics, kinetics and macroscopic mechanical behavior, and applications have been described (Delaey et al., 1974; Krishnan etal., 1974; Warlimont e t a l , 1974; Ahlers, 1986; Van Humbeek and Delaey, 1989; Hornbogen, 1991). 11.6.2 Cu-Zn-Al Shape Memory Alloys The Cu-Zn-Al shape memory alloys are derived from the B2 CuZn phase by alloying with Al, and the composition of the

737

resulting ternary phases is approximately described by Cu 1 + 1 3 x Zn 1 _ 2 ^Al^ (Van Humbeek and Delaey, 1989; Hodgson, 1990). The complete ternary Cu-Zn-Al phase diagram was given in Guertler et al. (1969). The crystal structure of this Cu-ZnAl phase is the disordered A2 structure at high temperatures which transforms to the B2 structure by nearest neighbors ordering on cooling and to the L2i structure by next-nearst neighbors ordering on further cooling (Ahlers, 1986; Wu and Wayman, 1991). It has to be emphasized that these structures are stable only at elevated temperatures whereas at room temperature they are only metastable, i.e. these alloys decompose if annealed at too low temperatures. The L2 t phase can be transformed martensitically to form various types of martensite which differ by the stacking sequence of the close-packed crystallographic planes (Ahlers, 1986). The crystallography and thermodynamic stability of the martensite variants have been analyzed in detail (Bidaux and Ahlers, 1992; Ahlers and Pelegrina, 1992; Pelegrina and Ahlers, 1992a, b; Saule et al., 1992). The martensite stability is correlated with elastic anomalies, as in the case of CuZn (Verlinden and Delaey, 1988 b). It can be shown that different types of martensite can coexist at certain compositions (Segui et al., 1991). In any case, the martensite formation temperature is a very sensitive function of the composition (Van Humbeek and Delaey, 1989; Hodgson, 1990), and can be varied by prior thermal cycling (Tadaki etal., 1987). In particular, alloying with Mn is used to depress the martensite formation temperature. As already mentioned for CuZn, the characteristics of such shape memory alloys have been discussed in detail with respect to crystallography, microstructure,

738

11 Intermetallics

thermodynamics, kinetics, and macroscopic mechanical behavior (Delaey et al., 1974; Krishnan etal., 1974; Warlimont e t a l , 1974; Ahlers, 1986; Van Humbeek and Delaey, 1989; Hornbogen, 1991). A special feature of the Cu-Zn-Al alloys is the so-called two-way memory effect which produces a reversible shape change on cooling and heating (Stalmans et al., 1992 a, b, c). This effect results from martensite transformation and retransformation after specific thermomechanical treatments which are designated as "training". The processing of Cu-Zn-Al alloys poses few problems (Hodgson, 1990). The alloys are usually produced by induction melting. Various dopants are in use for grain refinement to improve formability (Lee and Wayman, 1986). Apart from this, powder metallurgy is used for obtaining finegrained materials. The alloys are readily hot worked in air and the cold workability decreases with increasing Al content. Applications have been described (e.g. by Hodgson, 1990; Hornbogen, 1991). 11.6.3 Cu-Al-Ni Shape Memory Alloys The Cu-Al-Ni shape memory alloys are based on the intermetallic phase Cu3Al with a disordered A2 structure, which is usually known as the (3 phase, and which is stable only at high temperatures between 567 and 1049 °C with compositions between 71 and 82 at.% Cu (Murray, 1985). At 567 °C this phase decomposes by a eutectoid reaction at such a sluggish rate that it can be retained with the then metastable A2 structure by cooling below this temperature. At about 500 °C the p phase undergoes an ordering reaction to form the metastable p t phase with a D0 3 structure. Both phases can be transformed martensitically by quenching to form various types

of martensite depending on the Al content and the quench rate (Murray, 1985). Cu3Al can be alloyed with Ni without changing the transformation characteristics of this phase (see Guertler et al., 1969 for the ternary Cu-Al-Ni phase diagram). Both binary Cu3Al and ternary (Cu,Ni)3Al show the shape memory effect due to thermoelastic martensite formation (Otsuka and Shimizu, 1970; Nagasawa and Kawachi, 1971; Delaey etal., 1974). This ternary phase has given rise to the development of the Cu-Al-Ni shape memory alloys with about 11-15 wt.% Al and 3-5 wt.% Ni (Van Humbeek and Delaey, 1989; Hodgson, 1990). The crystallography of martensite formation has been discussed recently in detail for Cu-Al-Ni alloys (Shimizu and Tadaki, 1992), and likewise the thermodynamics and kinetics have been studied (Warlimont et al., 1974; Zhou and Hsu, 1991; Ortin etal., 1992). The deformation modes of the martensites include various twinning modes and stressinduced transformations (Ichinose et al., 1991). The shape memory effect is a twoway effect, i.e. it is completely reversible (Sakamoto etal., 1991). The martensite start temperature has an upper limit of about 180 °C and can be lowered to a large degree by increasing the Al content, and to a smaller degree by increasing the Ni content (Van Humbeek and Delaey, 1989; Hodgson, 1990). The Cu-Al-Ni alloys are advantageous because of their higher stability at higher temperatures compared with the Cu-Zn-Al alloys. However, second-phase precipitation cannot be suppressed and embrittles the Cu-Al-Ni alloys and precludes cold working, i.e. such alloys can only be hot finished (Van Humbeek and Delaey, 1989; Hodgson, 1990). Thermomechanical treatments and microalloying additions - in particular Mn, Ti, and Zr - are used for

11.6 Cu-Base Phases

improving the mechanical behavior. An important effect is grain refinement, which increases both the fracture strength and the fracture strain according to a HallPetch relationship (Roh et al, 1991). Indeed precipitated particles have been found to reduce the as-cast grain size and hinder later grain growth during annealing (Ratchev et al., 1993). The alloying additions affect the transformation and decomposition behavior which is to be considered for establishing the optimum conditions for the induction of the two-way effect (Eucken etal., 1991; Hornbogen and Kobus, 1992). Recently the ductilizing effect of small boron additions has been found, which has been studied in detail with respect to transformation behavior and microstructure (M. A. Morris, 1991 b, 1992). Finally it is noted that there are other similar Cu-base alloys which also show the shape memory effect (Delaey et al., 1974). A well-studied example is given by the Cu-Al-Mn shape memory alloys which order by a two-step reaction on rapid cooling, i.e. the initial, disordered A2 structure is transformed first to the intermediate B2 structure and then to the final L2X structure (Guilemany and Peregrin, 1992; Prado etal., 1993). 11.6.4 Cu-Au Phases The Cu-Au system is the classic textbook example for discussing ordering reactions in solid solutions and the effects of atomic order on properties (see, e.g., Schulze, 1967; Honeycombe, 1968). At higher temperatures above 410 °C the CuAu alloys form the disordered Al structure with complete mutual solid-solubility of Cu and Au, whereas at lower temperatures ordering reactions occur which produce various intermetallic phases, depending on temperature and composition, with broad

739

ranges of homogeneity (Okamoto et al., 1987). On the Cu-rich side there is the wellknown phase Cu3Al with an Ll 2 structure, which is the prototype phase for this crystal structure and which is designated Cu 3 Au I. The phase Cu 3 Au II forms at intermediate temperatures with off-stoichiometric compositions and is described as either tetragonal or orthorhombic. In any case, the Cu 3 Au II structure results from the L l 2 structure by the introduction of equally spaced antiphase boundaries (APBs). Such a structure is usually known as a long-period superlattice (LPS). At the equiatomic composition the tetragonal L l 0 phase CuAu I forms, and again an orthorhombic LPS is observed at intermediate temperatures, which is known as CuAu II. The Au-rich phase Au 3 Cu shows the LI 2 structure as Cu 3 Au. All phases are separated from each other by two-phase equilibria. The elastic behavior is known (Guillet and Le Roux, 1967) as well as the diffusion behavior (Wever et al., 1989; Wever, 1992) and vacancy formation (Schaefer et al., 1992). The deformation behavior of Cu 3 Au was studied in detail to clarify the effects of atomic order, and it was found in particular that order decreases the yield stress and increases the work hardening rate, with still high ductility (Sachs and Weerts, 1931; Vidoz etal., 1963; Stoloff and Davies, 1966). Inversely, strong cold deformation destroys the order, as was observed for various phases with ordering reactions (Dahl, 1936). The ordering reaction produces antiphase domains which give rise to an additional strengthening effect in this and similar phases, with strength increasing with decreasing domain size (Marcinkowski, 1974a). The domain size can be varied by prior processing, and it increases during ageing in analogy to Ostwald ripening (Sauthoff, 1973 a, b).

740

11 Intermetallics

The Portevin-Le Chatelier effect with serrated yielding was observed for both the ordered and disordered state (Mohamed et al., 1974). Recovery and recrystallization have been analyzed in detail (Vidoz etal., 1963; Cahn, 1990, 1991). Experimental and theoretical studies have been directed at dislocations and antiphase domain boundaries (see, e.g. Tichelaar and Schapink, 1991; D. G. Morris, 1992; Veyssiere, 1992), grain boundaries (Yan et al., 1992), and the electronic structure (Bose etal., 1991). It is noted that disordered layers are formed in ordered Cu 3 Au on antiphase boundaries and twin boundaries just below the order-disorder transition temperature (Tichelaar et al., 1992). This may be expected in other phases, too, and may improve the ductility of less ductile phases, as is discussed for Ni3Al (see Sec. 11.4.1.2). Au-Cu-Ag alloys based on the intermetallic phases CuAu and Cu 3 Au have found applications in dentistry because of their extremely high corrosion resistance, their advantageous mechanical properties such as high strength and ductility, and their decorative gold color (Yasuda, 1991). These alloys age-harden as a result of complex ordering and decomposition reactions by which the phases Cu 3 Au I, CuAu I, CuAu II, and an Ag-rich oc2 phase are formed, depending on the composition. 11.6.5 Cu Amalgams Dental amalgams have been in use for more than a thousand years (Waterstrat, 1990). Amalgams are still a very important class of material for dental restoration, in spite of corrosion problems with mercury release because the amalgams offer an advantageous combination of properties with respect to processing and use (Watts, 1992; Waterstrat and Okabe, 1994). These

amalgams are formed by alloying Ag-SnCu alloys based on the phases Ag3Sn and Cu3Sn, with mercury. The amalgamation reaction results in the formation of a complex multiphase alloy structure which contains the phases Ag 2 Hg 3 , Sn8Hg, Ag3Sn, and Cu 6 Sn 5 , depending on the composition. The processing with mixing and compaction of powders is comparatively simple and is analogous to mechanical alloying. The amalgams must show sufficient compressive strength and creep resistance since the oral temperature is only about 10% below the melting temperature. The dental amalgams are very brittle, which leads to failure with tensile or bending stresses in unsupported regions. The fracture processes are correlated with corrosion processes which primarily attack the phases Sn8Hg and Cu 6 Sn 5 and lead to the release of mercury and SnO, respectively. Attempts to substitute for mercury in these amalgams have not yet been successful.

11.7 A15 Phases 11.7.1 Basic Properties The A15 crystal structure is a complex, cubic structure derived from the b.c.c. lattice (A2) with 8 atoms per unit cell (Pearson, 1958). It is a typical example of the topologically close-packed (tcp) structures which result from the stacking of polyhedra of various shapes to accommodate atoms of different sizes (Sinha, 1973; Watson and Bennet, 1984; Gladyshevskii and Bodak, 1994). These tcp structures allow a closer packing of atoms of different sizes than the geometrically close-packed f.c.c. (Al) and h.c.p. (A3) structures. The A15 structure may be regarded as the basis of the family of tcp structures since the other

11.7 A15 Phases

known tcp structures - in particular the Laves phase structures C14, C15, and C36 (see Sec. 11.8) - can be derived from the A15 structure by various crystallographic operations (Ye etal., 1985). The intermetallic phases with tcp structures are known as Frank-Kaspar phases. The phases with A15 structures form a large and comparatively homogeneous group of intermetallic phases (Nevitt, 1967) which have been the subject of fundamental studies, e.g. with respect to electronic structure and phase stability (Turchi and Finel, 1991) or the micromechanisms of diffusion (Bakker and Westerveld, 1988; Bakker etal., 1985, 1992; Wever etal., 1989; Koiwa, 1992; Wever, 1992), twinning (Khantha etal., 1989), and plastic deformation (chapter by Umakoshi in Volume 6 of MST). The A15 phases are of enormous practical importance since many of them are superconductors (B. W. Roberts, 1967; Geballe and Hulm, 1975; Furuto, 1984; Dew-Hughes, 1986; Muller, 1986; Stekly and Gregory, 1994). Some of them, i.e. V 3 Ga, V3Si, Nb 3 Al, Nb 3 Ga, Nb 3 Ge, and Nb 3 Sn, exhibit critical temperatures between 15 K and 23 K which are only surpassed by those of the ceramic superconductors based on the perovskite-type oxides (Bednorz and Muller, 1988; DewHughes, 1988; Sharp, 1991); see also the chapter by Clarke and Daumling in Volume 11 of this Series. However, the latter are inferior with respect to critical current density, i.e. the best combination of critical temperatures, critical current density and critical magnetic field strength is found in the A15 superconductors compared with all other superconducting materials. Recently A15 phases - in particular Nb 3 Al and Cr3Si - have also been considered for applications as structural materials at high temperatures (Shah et al., 1990; Shah and Anton, 1992a; Kamata etal.,

741

1992 a, 1993; Suyama et al., 1993). The application of A15 phases as structural or functional materials is handicapped by the brittleness of the A15 phases, which makes processing difficult. In the following, a brief overview is given of those A15 phases which have already been applied, or have been considered for application, and which have been studied in some detail. 11.7.2 V3Si V3Si was the first A15 phase discovered to be superconducting, with a critical temperature of 17K (Smathers, 1990). Stoichiometric V3Si melts congruently at 1925 °C and is stable at high and low temperatures (Massalski et al., 1990). A deviation from stoichiometry is possible on the Si-poor side which, however, reduces the critical temperature due to the resulting constitutional disorder. The critical temperature is also reduced by dissolved hydrogen (Stepanov and Skripov, 1982). Likewise the critical current density and the critical field strength are reduced by elastic tensile or compressive straining (Smathers, 1990). Slightly above the critical temperature V3Si transforms martensitically to a weakly tetragonal, still superconducting phase with twin-related thin lamellae, which is due to lattice softening with vanishing shear modulus (Batterman and Barrett, 1964, 1966; Kondratyev and Pushin, 1985; Mendelson, 1986; Onozuka etal., 1988). Prior plastic deformation with large strains at high temperatures stabilizes the tetragonal martensite even at temperatures above the martensite formation temperature (Ullrich etal., 1978). The martensitic transformation gives rise to superelasticity and a shape memory effect (Zakrevskiy etal., 1986). V3Si is hard and brittle, as are all the A15 phases, which makes processing very

742

11 Intermetallics

difficult (Fleischer, 1989; Fleischer and Zabala, 1990 c). The reason for this is the electronic structure of V3Si which leads to strong covalent bonding between the nearest neighbor V atoms (Medvedev et al., 1984; Muller, 1986). V3Si has a high Young's modulus of 213 GPa (Fleischer e t a l , 1989b) and an apparent brittle-toductile transition temperature (BDTT) of the order of 1200 °C (Mahajan et al., 1978; Nghiep et al., 1980). The plastic deformation behavior above the BDTT has been studied in detail and the dislocation slip systems have been analyzed (Levinstein etal., 1966; Bertram and Paufler, 1983; Kramer, 1983; Wright and Bok, 1988; Smith etal., 1993). The creep behavior is characterized by subgrain formation, dislocation climb controlling the creep rate and a comparatively high activation energy in the range of 2-11 eV/atom depending on the stress and the composition (Nghiep et al., 1980). Creep can be enhanced by superimposed electric currents, i.e. there is the so-called electroplastic effect, and thermal gradients, which is proposed to be due to the favorable effect of the electrotransport and thermotransport of point defects on the dislocation mobility (San Martin et al., 1980, 1983). The transition from superconduction to normal conduction in A15 superconductors occurs by the penetration of small fluxoids, i.e. quantized magnetic flux vortices, which is characteristic for the type II, high-field superconductors (Geballe and Hulm, 1986). These fluxoids can be pinned by microstructural inhomogeneities - in particular grain boundaries and precipitated particles - which is beneficial with respect to critical current density and field strength (Muller, 1986; Smathers, 1990). Consequently, microstructural changes by prior plastic straining with or without recrystallization affect the superconducting

properties, depending on the deformation conditions and the composition (Quyen et al., 1979). The thin lamellae, which are produced by the already-mentioned martensitic transformation, are less effective with respect to flux pinning than grain boundaries, i.e. their effect on the superconducting properties becomes relevant only in single crystals (Brand and Webb, 1969). V3Si is of interest for applications because of its favorable superconducting properties and its high stability (Smathers, 1990). Apart from hot extrusion, V3Si can be deformed plastically at room temperature in spite of its brittleness by superimposing a high hydrostatic pressure which makes the production of multifilament, superconducting wires feasible for applications in solenoid magnets (Wright, 1977). Laminar V3Si can be prepared by solidstate, thin film reactions between V and SiO2 on Si substrates (Hayashi et al., 1991). However, Nb 3 Sn has been found to be more favorable with respect to superconducting properties, mechanical behavior, and processing, which has precluded the use of V 3 Si as a magnet material (Smathers, 1990). Recently V3Si has become of interest with respect to structural, high-temperature applications (Shah and Anton, 1992 a). Combination with a ductile phase, e.g. Vrich solid-solution, offers the possibility of overcoming the brittleness problem, and indeed a directionally-solidified, eutectic V 3 Si-V alloy has been successfully prepared (Goldman, 1993). The fracture toughness of V 3 Si-V composites with discontinuous V3Si has been found to increase monotonically with increasing fraction of V and decreasing impurity content (Strum and Henshall, 1993).

11.7 A15 Phases

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11.7.3 V3Ga

11.7.4 Nb3Sn

V 3 Ga with an A15 structure is stable only up to 1300°C where it transforms congruently to a V-Ga disordered, b.c.c. solid-solution (Massalski etal., 1990). There are indications of a martensitic transformation at very low temperatures, in close analogy to V3Si, though the conditions for martensite formation are not yet clear for V 3 Ga (Nembach etal., 1970; Snead, Jr. and Bussiere, 1985). It has a broad range of homogeneity centered on the stoichiometric composition, i.e. it is very stable with respect to deviations from stoichiometry. Because of this stability, which is advantageous for fabrication, and the still high, critical temperature of about 16 K, V 3 Ga is of interest for applications (Smathers, 1990). The superconducting properties and the deformation behavior of V 3 Ga are similar to those of V3Si. The superconducting transition can be affected by prior plastic deformation (Wright and Ho, 1986). V 3 Ga is ductile in compression above 1000 °C and the creep behavior has been studied in detail (Soscia and Wright, 1986; Shah and Anton, 1992 a). As in the case of V3Si, V 3 Ga can be deformed plastically by hot extrusion and at room temperature by superimposing a high hydrostatic pressure, which makes the production of multifilament, superconducting wires feasible for applications in solenoid magnets (Wright, 1977). Hybrid high-field magnets with V 3 Ga tape coils have been produced successfully (Noto etal., 1986). However, Nb 3 Sn has been found to be more favorable with respect to superconducting properties, mechanical behavior, and processing which has precluded the commercialization of V 3 Ga up to now (Smathers, 1990).

Nb 3 Sn with an A15 structure forms peritectically at 2130°C from the liquid solution and solid Nb, and is stable only above 775 °C where it decomposes eutectically (Massalski etal., 1990). The decomposition reaction can be suppressed by not too slow cooling, i.e. Nb 3 Sn with an A15 structure is maintained at lower temperatures in a metastable condition. Nb 3 Sn transforms to a twinned, tetragonal martensite at 43 K (Mailfert etal., 1969; Mendelson, 1986), and the transition to superconductivity occurs at about 18 K (Muller, 1986; Smathers, 1990). This critical temperature is reduced by deviations from stoichiometry, which affect the atomic order, and by elastic strains. Small increases in the critical temperature are possible using ternary alloying additions. The transition temperature does not depend sensitively on prior plastic deformation (Wright and Ho, 1986). As for the other A15 phases, Nb 3 Sn can be deformed readily at high temperatures, e.g. by hot extrusion, whereas deformation at low temperatures is possible only by superimposing a high hydrostatic pressure (Wright, 1977; Eisenstatt and Wright, 1980; Clark and Wright, 1983). The high temperature deformation and creep behavior have been studied in detail and a beneficial effect of grain size reduction on the brittle-to-ductile transition temperature (BDTT) has been found (Clark et al., 1983; Shah and Anton, 1992 a). Nb 3 Sn superconductors with multifilamentary geometry, which is needed for electromagnet applications, are easily fabricated by the so-called bronze process, which has led to the wide-spread commercial application of Nb 3 Sn (Smathers, 1990). In the bronze process Nb rods are inserted in a bronze rod or tube which is

744

11 Intermetallics

then processed to obtain a multifilament wire. At the end of the wire processing, Nb 3 Sn shells are formed around the Nb cores by reaction heat treatments. The final multifilament, composite superconductor consists of Nb 3 Sn fibers with Nb cores which are embedded in a bronze matrix; the latter is surrounded by a coating of Nb, which serves as a diffusion barrier, and an outer Cu tube, which serves as a stabilizer and carries the current in case of transition to normal conduction. Various specific processing routes, including powder metallurgy routes, have been developed for optimized superconductor production and hybrid superconductor coils are commercially available (Hillman et al, 1980; McDonald, 1984; Gregory, 1984; Noto etal., 1986; Hecker etal., 1988; Smathers, 1990). The superconducting properties of the composite superconductor, as well as the mechanical properties, depend on the particular microstructure which is a function of the various prior processing steps. Much work has been directed towards the details of the various processes, e.g. the effects of alloying additions (Zwicker etal., 1979), the Sn diffusion during Nb 3 Sn formation (Glowacki and Evetts, 1988; Cheng and Verhoeven, 1988), the grain size dependence of the flux pinning force and the tensile strength and fracture at room temperature (Ochiai et al., 1986a, b, 1988; Watari et al., 1986), or the effects of the hardnesses of the constituent phases on the composite workability (Dew-Hughes etal., 1987). Recent new developments rely on the beneficial effects of additions of Ti, Hf, Ta, and/or Ge on the superconducting properties (Kohno et al., 1992; Kamata et al., 1992b; Tachikawa et al., 1992; Murase et al., 1992; Noto et al., 1992).

11.7.5 Nb3Al Nb 3 Al with an A15 structure forms peritectically at 2060 °C and is stable below this temperature in the whole temperature range only with Al-deficient, off-stoichiometric compositions (Jorda etal., 1980; Massalski et al., 1990). The stoichiometric composition corresponds to the maximum solubility of Al in Nb 3 Al, and is reached only at 1940 °C in eutectic equilibrium with the liquid Nb-Al and the a phase Nb 2 Al. The phase stability is a function of the electronic band structure which has been studied as a function of temperature (Kuzmichev etal., 1983). The formation enthalpy of Nb 3 Al is significantly smaller than that of the Al-rich phase Al 3 Nb, which has been discussed in Sec. 11.3.3 (Meschel and Kleppa, 1993). The transition to superconduction occurs at 19 K and there is no martensitic transformation (Smathers, 1990). The transition temperature of Nb 3 Al is affected by deviations from stoichiometry and perfect order, as are those of the other A15 phases, but in a less sensitive way. It is noted that the A15 phase with the highest transition temperature is Nb 3 Ge, which can be alloyed with Nb 3 Al to form the ternary A15 phase Nb 3 (Al 0 7 Ge 0 3) with the highest critical field strength (Muller, 1986). Nb 3 Al has long been considered for electromagnet applications because of its advantageous superconducting properties (Smathers, 1990). However, the processing of Nb 3 Al is very difficult because of the brittleness of Nb 3 Al. Powder metallurgy methods have been used successfully for preparing Nb 3 Al by reaction synthesis and producing, e.g. an Nb 3 Al-Ag composite superconductor (Bowden, 1989; Tsuchida etal., 1989; Schulze etal., 1990). Nb 3 Al tapes have been prepared by rapid solidification (Togano et al., 1992). A novel pro-

11.7 A15 Phases

cess with thin Nb-Al sandwiches as the starting material has allowed the fabrication of Nb 3 Al superconducting wires with good workability (Saito et al., 1990, 1993; Ikeda et al., 1992). Furthermore, various other processes have been used successfully for producing Nb 3 Al wires (Noto et al., 1992). Presently Nb 3 Al is also being considered for structural applications at high temperatures in spite of its low-temperature brittleness because of its high stability and strength at high temperatures, and various developments have been initiated (Dimiduk et al., 1991; Anton and Shah, 1991a; Yamaguchi, 1992; Bunk, 1992; Kamata et al., 1993; Suyama et al., 1993). The mechanical properties at low and high temperatures have been determined and the brittle-to-ductile transition has been observed at about 1000°C (Barth et al., 1992; Shah and Anton, 1992 a). At lower temperatures Nb 3 Al is inherently brittle according to theoretical ab initio calculations (Kim et al., 1993). The dislocation configurations and slip systems have been analyzed (Aindow et al., 1991; Murayama etal., 1993 a-c). The yield stress can be varied by varying the Al content and by alloying with a third element (Fujiwara etal., 1991; Kamata etal., 1992a). The solubilities of third metals in Nb 3 Al differ enormously depending on the particular metal, which offers manifold possibilities of alloying (Shah and Anton, 1992 a). The diffusion behavior has been studied (Slama and Vignes, 1972), and the oxidation resistance is only poor because of nonprotective scale formation (Shah and Anton 1992a; Tomizuka, 1992; Fujiwara etal., 1993). A reduction in the Al content below the solubility limit leads to two-phase Nb 3 AlNb alloys which offer the possibility of relieving the brittleness of Nb 3 Al (Anton

745

and Shah, 1990; Marieb etal., 1991; Kumagai etal., 1991; Davidson and Anton, 1993). The brittle-to-ductile transition temperature (BDTT) decreases with increasing volume fraction of the Nb-rich phase, and a brittle-to-ductile transition at room temperature has been obtained for an Nb-16 at.% Al alloy (Suyama and Hashimoto, 1992). The Nb phase toughens the alloy by crack bridging, plastic stretching, and interfacial bonding (Murugesh et al., 1992). However, the fracture toughness controlling, microstructural details are not yet clear (Davidson and Anton, 1993). It is noted that the Nb-rich phase, which usually crystallizes with the b.c.c. A2 structure, may be ordered with a B2 structure due to the presence of interstitial impurities (Marieb et al., 1991). The transformation and precipitation reactions have been studied in detail (Yang and Vasudevan, 1993). Powder metallurgy methods and infiltration techniques with reaction synthesis result in such two-phase alloys with even higher Al contents because of incomplete reactions (Kumagai et al., 1991; Murayama et al., 1991). Nb-Al alloys with still higher Al contents in the range 25-33 at.% contain the a phase Nb 2 Al, which is also regarded as attractive for high-temperature applications (Bhattacharya et al., 1992). Apart from the above Nb 3 Al-Nb alloys, which may be regarded as discontinuously reinforced, in situ composites, Nb 3 Al has been considered for use in intermetallic matrix composites which are reinforced by continuous ceramic fibers (Anton, 1988; Shah et al., 1990). Laminated composites have been prepared in situ by high rate sputtering (Rowe and Skelly, 1992). Alloying of Nb 3 Al with Ti leads to a ternary phase with a B2 structure (Shyue et al., 1993) which is related to the Ti3Albase, super-oc2 alloys of Sec. 11.3.1.2.

746

11 Intermetallics

11.7.6 Nb3Si Nb3Si is a high-temperature line compound which forms with a tetragonal crystal structure by a peritectic reaction at 1980 °C and is stable only above 1770°C (Massalski etal., 1990). Rapid solidification of the liquid produces amorphous Nb 3 Si (Bertero etal., 1991a). The A15 structure is unstable and is obtained only by special rapid solidification techniques and by crystallization of amorphous Nb 3 Si. Nb 3 Si with an A15 structure is of interest because a high transition temperature for superconduction is expected, and indeed a transition temperature of 19 K was reported for near-stoichiometric Nb 3 Si with an A15 structure, which was prepared by crystallization of amorphous Nb 3 Si under high pressure (Wang et al., 1988). Other crystal structures are obtained for Nb 3 Si doped with impurities, i.e. the L l 2 structure with oxygen and a hexagonal structure with carbon (Kassem and Koch, 1991). The present interest is directed at Nb-Si in situ composite alloys which contain the Nb 3 Si phase and are considered for high-temperature applications (Bertero etal., 1991a, b; Cockeram etal., 1991, 1992a, b; Bewlay etal., 1992; Goldman, 1993). 11.7.7 Cr3Si Cr3Si crystallizes congruently with an A15 structure from the melt and is stable down to very low temperatures with a range of homogeneity which includes the stoichiometric composition (Massalski et al., 1990). Cr3Si does not exhibit superconduction (Smathers, 1990). However, Cr3Si has been regarded for a long time as promising for high-temperature applications because of its high creep strength and oxidation resistance (Arbiter, 1953 a, b; Silverman, 1956; Anton and Shah, 1991b;

Dimiduk etal., 1991). The elastic moduli have been determined as well as the temperature dependence of the hardness (Fleischer etal., 1989b, 1991; Fleischer and Zabala, 1990 c). Compressive ductility has been observed only above 1200 °C and its extent is expected to be limited because of an insufficient number of slip systems (Chang and Pope, 1991). Creep data are available (Shah and Anton, 1991). At lower temperatures Cr3Si is brittle with a low fracture toughness (Fleischer, 1990; Fleischer etal., 1991). Sufficient oxidation resistance has been observed only below 1200 °C, i.e. below the apparent brittle-toductile transition temperature (McKee and Fleischer, 1991). An improved toughness may be obtained with two-phase Cr 3 Si-Cr alloys, i.e. by combining the brittle Cr3Si phase with the Cr-rich solid solution (Mazdiyasni and Miracle, 1990; Newkirk and Sago, 1990; Bewlay etal., 1992). An improved mechanical behavior is also expected for Cr 3 Si-based, intermetallic matrix composites, e.g. Cr 3 SiTiC, Cr 3 Si-TiB 2 , or Cr 3 Si-Y 2 O 3 , for which good bonding and chemical compatibility of matrix and reinforcements have been found (Yang et al., 1990).

11.8 Laves Phases 11.8.1 Basic Properties The Laves phases - sometimes designated as Friauf-Laves phases - with an AB 2 composition in the binary case form a very large group of intermetallics which crystallize with the hexagonal C14 structure, the cubic C15 structure, or the dihexagonal C36 structure (Laves, 1967; Wernick, 1967; Livingston, 1992). These structures are topologically close-packed (tcp) structures (Wernick, 1967; Schulze etal., 1973; Watson and Bennet, 1984), i.e. the Laves

11.8 Laves Phases

phases belong to the family of the FrankKaspar phases, which has already been discussed in Sec. 11.7 in the context of the A15 phases. The stability of the three crystal structures C14, C15, and C36 is controlled by both the atomic size ratio of the A atoms and B atoms and by the valence electron concentration of the Laves phase (Wernick, 1967; Leitner, 1971a, b; Leitner and Schulze, 1971; Ohta and Pettifor, 1990). Deviations from stoichiometry are accommodated primarily by placing the excess atoms on sites of the other sublattice, i.e. the A and B atoms substitute for each other (Bruckner, 1969; Schulze, 1972). Such composition changes affect the valence electron concentration which may induce a change in the crystal structure, as has been observed in the case of TiCo 2 which shows the C15 structure for the stoichiometric composition and the C36 structure for the Co-rich, off-stoichiometric composition (Aoki et al., 1966; Schulze, 1972). It is noted that the respective phase equilibria are still in discussion (Massalski et al., 1990). A similar situation has been reported for TaCo2 with the additional C14 structure for the Co-poor, off-stoichiometric composition and for NbCo 2 (Massalski et al., 1990). Different crystal structures have also been observed at different temperatures for some Laves phases, e.g. NbCo 2 and the chromides MCr 2 with M = Ti, Zr, Nb, or Ta, which indicates only small energy differences between the three crystal structures (Massalski e t a l , 1990). Indeed the crystal structures C14, C15, and C36 differ only by the particular stacking of the same two-layered structural units which allows structure transformations between these structures and twinning by synchroshear (Allen et al., 1972; Allen and Liao, 1982; Allen, 1985; Hazzledine et al., 1993; Hazzledine and Pirouz, 1993; Liu et al., 1993;

747

Pope and Chu, 1993). Correspondingly, stacking sequence faults may be formed easily in topologically close-packed phases, which have been studied in detail (Khantha etal., 1990). A structure transformation may be induced by a mechanical stress, a has been observed for ZrFe 2 with a C36/C15 transition (Liu etal., 1993). The basic properties have only been determined for a few selected Laves phases, e.g. elastic moduli (Shannette and Smith, 1969; Schulze and Paufler, 1972; Schiltz and Smith, 1974; Balankin, 1984; Halstead and Rawlings, 1985; Fleischer etal. 1988; Fleischer and Zabala, 1990 b, c; Fleischer, 1992 b) (see also Fig. 11-6), thermal expansion (Giegengack et al., 1966), and diffusion properties (van der Straten et al., 1976; Wein etal., 1978; Wever, 1992; Sprengel etal., 1994). The elastic anisotropy is only small which indicates the absence of strong directional bonding (Schulze and Paufler, 1972). HfV2 exhibits an anomalous temperature dependence of the elastic moduli which is related to a martensitic transformation at low temperatures (Livingston and Hall, 1990). The plastic deformation behavior at high and low temperatures, i.e. dislocation slip systems and mobilities, has been studied primarily for the classic Laves phases MgZn 2 with a C14 structure and MgCu 2 with a C15 structure (Moran, 1965; Paufler and Schulze, 1967; Kramer and Schulze, 1968; Paufler, 1972; Hall and Livingston, 1989; Livingston et al. 1989; Ohba and Sakuma, 1989). Recent studies have been directed at HfV2 with a C15 structure (Hall and Livingston, 1989; Livingston and Hall, 1990; Pope and Chu, 1993), TiCr2 with a C15 structure (K. C. Chen et al., 1993), MgNi 2 with a C36 structure (Livingston and Hall, 1991), and the ternaries Mg(Cu,Zn)2 with a C36 structure (Livingston and Hall, 1991) and Nb(Ni,Al)2 with a C14 structure

748

11 Intermetallics

(Sauthoff, 1990 b, 1991a; Machon, 1992). The Laves phases are strong and brittle with an apparent brittle-to-ductile transition temperature of about two thirds of the melting temperature (Wetzig and Wittig, 1972; Schulze and Paufler, 1972; Livingston, 1992). This brittleness is also shown by single crystals, i.e. it is not due to the presence of weak grain boundaries and an insufficient number of slip systems. It is an inherent brittleness which results from the complexity of the dislocation glide processes with high Peierls stresses and correspondingly low dislocation densities and mobilities. The already high strengths of Laves phases can be further increased by solidsolution strengthening as a result of alloying with third elements (Sauthoff, 1989; Livingston, 1992; Fleischer, 1992 b, 1993 d). The alloying elements can substitute for both the A atom or the B atom in the binary AB 2 phase, as is indicated by the extension of the AB 2 field in the isothermal sections of the respective ternary phase diagrams (Anton and Shah, 1991 b). Alloying with Al and Si is of particular interest in view of the beneficial effect on the oxidation resistance (Meier and Pettit, 1992). Alloying with both elements, which substitute for the B element in AB 2 , may change the valence electron concentration to such an extent that a change in crystal structure is induced, i.e. each of the three alternative crystal structures C14, C15 and C36 can be stabilized by the appropriate control of the Al and Si contents (Bardos e t a l , 1963; Wernick, 1967; Wallace and Craig, 1967). A similar situation has been found for the ternary Nb(Cr,Fe) 2 which forms with the C14, C15 or C36 structure depending on temperature and Fe content (Grujicic et al., 1993). The effect of both composition and temperature on structure stability is exem-

plified by the ternary phase Nb(Cox _ ^ A l ^ which exhibits the C14 structure at high temperatures and the C15 structure at low temperatures, with a transition temperature which decreases with increasing Al content from about 1300°C for x = 0 to room temperature for about x = 0.12 (Von Keitz and Sauthoff, 1992). A similar behavior has been observed for Nb(Co1_JCSi)2. However, Ta(Fe 1 _ : c ,Ay 2 , Nb(Fe 1 _ x ,Al JC ) 2 , Nb(Fe x _ „ Six)2, Ti(Fex _ „ Al x ) 2 , and Ti(Fe1_JC,Six)2 only show the C14 structure (Raghavan, 1987). It has to be emphasized that some other similar ternary Laves phases, e.g. Ta(Ni1_x,AlJC)2, Nb(Ni 1 _ x ,AlJ 2 , and Mo(Co 1 _ x ,SiJ 2 all with C14 structures do not exist as binary phases, i.e. the cases x = 0 are not included and correspondingly these ternary Laves phases are frequently given as, e.g. TaNiAl, NbNiAl, and MoCoSi (Bardos et al., 1961; Benjamin et al., 1966; Skolozdra etal., 1966b; Villars and Calvert, 1991). Finally, it is noted that the alreadymentioned binary TiCr2 dissolves Fe to form the ternary Ti(Fe 1 _ x ,Cr x ) 2 with a C14 structure for 0 < x < 0.9, whereas the structure of this ternary phase is either C14, C15, or C36 for 0.95 < x < 1 depending on the composition and the temperature (Raghavan, 1987). 11.8.2 Applications

Various Laves phases have been regarded as promising for both functional and structural applications (Livingston, 1992), and examples are given in the following sections. 11.8.2.1 Superconducting Materials

A transition to superconductivity at low temperatures has been observed for a large number of Laves phases with the C14 structure or the C15 structure (B. W

11.8 Laves Phases

Roberts, 1967). (Hf,Zr)V2 with a C15 structure is of particular interest because of its advantageous superconducting properties since it combines a fairly high critical temperature with a high critical current density and magnetic field strength (Noto et al, 1986; Olzi et al., 1988). Processing is difficult because of its poor deformability. Further alloying - in particular with Nb improves the mechanical behavior to such an extent that multifilamentary superconductors can be produced which are being considered for application in fusion reactor magnets since the superconducting properties are comparatively insensitive to neutron irradiation and mechanical strain (Noto et al., 1986). Room temperature deformation of HfV2-base alloys occurs primarily by twinning (Hall and Livingstone, 1989; Livingston and Hall, 1990; Pope and Chu, 1993). The Hf-V-Nb phase diagram has been studied recently (Chu and Pope, 1992, 1993). 11.8.2.2 Magnetic Materials Various transition-metal Laves phases, e.g. TFe 2 with T = Ti, Zr, Hf, Nb, or Mo as well as NbCo 2 and the closely related rareearth Laves phases, show interesting magnetic properties as a result of their particular electronic structures (see, e.g. Bruckner, 1969; Buschow, 1980; Armitage e t a l , 1986; Yamada and Shimizu, 1986; Smirnova and Meshkov, 1986; Asano and Ishida, 1988). The rare-earth phases TbFe 2 and SmFe2, both with the C15 structure, exhibit huge magnetostrictions at room temperature (Clark, 1980). In addition there is a magnetoelastic interaction, i.e. unusual changes in the elastic moduli are observed under a magnetic field. The ternary C15 phase (Tb1_xDy;c)Fe2 has a high potential for application as a highpower transducer, since it combines a large

749

magnetostriction with a vanishing magnetic anisotropy (with x « 0.7) (Clark, 1980). This material is brittle like other Laves phases, whith an apparent brittle-to-ductile transition temperature of about 875 °C (Saka etal., 1991). Recently CeAl2 and TbAl 2 , both with C15 structures, have been studied with respect to the effect of microstructural defects on the magnetic behavior (Bi and Abell, 1993). Various rare earth (R) Laves phases, in particular RFe 2 , can absorb large amounts of hydrogen (see next section) which may affect their structures, i.e. they can undergo structure transformations on hydrogenation, including amorphization, which may be accompanied by decomposition reactions (Aoki and Masumoto, 1988; Pontonnier et al, 1991; Suzuki and Lin, 1993; Christodoulou and Takeshita, 1993 b; Kim and Lee, 1993 a, b; Aoki and Masumoto, 1993). Such hydrogenation and dehydrogenation reactions have been used for producing fine-grained powders for the powder metallurgical production of so-called giant magnetostrictive RFe 2 alloys (Jones etal., 1991). 11.8.2.3 Hydrogen Storage Materials Various Laves phases can absorb large amounts of hydrogen and are considered for applications as hydrogen storage materials (Somenkov and Shilstein, 1979; Reilly, 1979; Ivey and Northwood, 1986 b). Such AB 2 phases contain a strong hydride former as the A element and crystallize with the C14 structure or the C15 structure. The Laves phases with the highest sorption capacities are C15 ZrV2 with a hydrogen-to-metal ratio H/M = 1.8, C14 or C15 ZrCr 2 with H/M = 1.3, and C14 ZrMn 2 as well as C14 or C15 TiCr2 with H/M = 1.2. LaNi 2 , which was known as a C15 phase (Wernick, 1967), with

750

11 Intermetallics

H/M = 1.5 (Ivey and Northwood, 1986 b) is a tetragonal, pseudo-Laves phase which is stabilized only by the presence of impurities according to recent findings (Gschneidner, 1993). The formed hydrides in these binary Laves phases are too stable for easy hydrogen desorption, which is a prerequisite for hydrogen storage (Ivey and Northwood, 1986b). The hydride stability can be reduced and adjusted to practical hydrogen storage conditions by deviations from stoichiometry, substitution of the B element primarily by Fe, Co, Ni, Cu, Mn, or Cr, or substitution of the A element primarily by Ti, or any combination of these alloying possibilities. An example of the stoichiometry effect is given by ZrMn 2 , which exhibits an increasing hydrogen vapor pressure with increasing Mn content, and which is stable with a C14 structure up to ZrMn 3 8 (Ivey and Northwood, 1986 b). Another example is C14 TiMn 2 , which reaches its maximum H solubility at the off-stoichiometric composition corresponding to TiMnx 5 , whereas it does not absorb H at the stoichiometric composition (Sicking et al., 1981). The effects of substitution on electronic structure, alloy stability, hydride stability, and H storage capacity have been studied with respect to various systems, e.g. (ZrJCTi1_x)Mn2 (Moriwaki et al., 1991a), Zr(Mn 1 _ x Fe x ) 2+y (Triantafillidis etal., 1991), and Zr(V1_xCo3C)2 where both V and Co have been substituted for by other transition metals (Peretz et al., 1979), and Zr(Cr 1 _ x V x ) 2 (Ivey and Northwood, 1986a). Zr(Cr 1 _ x VJ 2 exhibits polytypism, i.e. crystal structure variants of either the hexagonal C14 structure or the cubic C15 structure are formed which only differ by the stacking of the crystal structure units (Meng and Northwood, 1986; Burany and Northwood, 1991). Var-

ious rare earth (R) Laves phases undergo amorphization on hydrogenation (Aoki and Masumoto, 1988, 1993; Pontonnier etal., 1991; Suzuki and Lin, 1993; Christodoulou and Takeshita, 1993 b; Kim and Lee, 1993 a, b). RNi 2 with a C15 structure is easily transformed into amorphous RNi 2 which is very stable and is being considered for use as long-life electrodes in electrochemical nickel-hydrogen batteries (Miyamura etal., 1991). The thermodynamics of hydrogenation have been studied and hydride formation enthalpies have been determined (see, e.g. Lynch etal., 1979; Sicking etal., 1981; Ivey and Nothwood, 1986b; Uchida et al., 1986; Perevesenzew etal., 1988; Drasner and Balzina, 1991; Zeng et al., 1993; Luck and Wang, 1993; Park etal., 1993). The optimization of properties has led to multinary Laves phases, e.g., Z r - M n - C o - V alloys (Yonezu etal., 1991), Z r - M n - N i - V alloys (Wakao et al., 1991), Z r - M n - N i V-Cr alloys (Moriwaki etal., 1991b), or Z r - T i - M n - F e and Z r - T i - V - F e alloys (Park and Lee, 1992), which have been studied in detail in view of applications as hydrogen storage materials and electrodes in electrochemical nickel-hydrogen batteries. 11.8.2.4 Structural Alloys The ternary Laves phase MoCoSi, i.e. Mo(Co,Si)2 with a hexagonal C14 structure, has given rise to the development of wear resistant Co-Mo-Cr-Si alloys which contain large volume fractions of the Laves phase in a coarse distribution together with a Co-rich phase, and which are known as Tribaloys (Schmidt and Ferris, 1975; Halstead and Rawlings, 1984, 1985). Apart from their high strength and hardness they have an adequate fracture toughness of the order of 20 MN/m 3/2 .

11.8 Laves Phases

These advantageous mechanical properties result in an excellent wear resistance which - in combination with an excellent corrosion resistance in various environments - makes the Tribaloys very appealing for applications in a broad range of temperatures and environments. The strong binary Laves phases TaFe 2 and NbFe 2 have been considered for the development of Laves phase strengthened, ferritic Fe-base alloys for applications at elevated temperatures (Bhandarkar et al., 1976; Wert et al., 1979). Monolithic transition metal Laves phases - in particular TiCr2 - have long been regarded as promising for high-temperature applications because of their high strengths at high temperatures and sufficient oxidation resistances (Arbiter, 1953 a, b; Silverman, 1956; Grinthal, 1956, 1958). However, the pronounced brittleness of such Laves phases makes processing very difficult and has precluded any application of monolithic Laves phases as structural materials. As in other alloy systems, the combination of such Laves phases with ductile phases can reduce the brittleness to a tolerable level, and indeed two-phase Ti-TiCr 2 alloys, as well as Ti-Nb-(Ti,Nb)Cr 2 alloys, have shown good prospects for obtaining high strengths with acceptable room-temperature toughness (Fleischer and Zabala, 1990b; K. C. Chen etal., 1993). The Laves phase is present in these alloys as strengthening particles which are cracked during deformation with the cracks being stopped at the phase boundaries. Similar alloy systems with Laves phases have been screened extensively with respect to strength, brittle-toductile transition temperature and oxidation (Anton and Shah, 1992 a, 1993; Shah and Anton, 1992 b). The Nb-Cr system is regarded as a candidate for high-temperature applications because the NbCr 2 Laves

751

phase has a high congruent melting point of 1770°C, high strength and creep resistance, comparatively low density, potential oxidation resistance, and a wide solubility range which offers manifold possibilities for alloying (Svedberg, 1976; Thoma and Perepezko, 1990; Anton and Shah, 1991 b; Vignoul et al., 1991; Takeyama and Liu, 1991). Toughness is improved not only by embedding NbCr 2 as particles in a Cr matrix, but also by Cr or Nb particles in an NbCr 2 matrix (Anton and Shah, 1990; Takeyama and Liu, 1991). Similar toughening effects have been observed for Cr-ZrCr 2 and Cr-HfCr 2 (Mazdiyasni and Miracle, 1990). Apart from these twophase metallic/intermetallic alloys, which may be regarded as in situ composites, such Laves phases are also considered as matrices in intermetallic matrix composites with strengthening dispersoids and fibers (Shah et al., 1990; Yang et al., 1990). Cr-containing Laves phases may be less advantageous at very high temperatures above 1000 °C with respect to oxidation resistance because of the volatility of the Cr oxides (Hindam and Whittle, 1982). For applications at such high temperatures Al-containing phases are preferred, as already remarked upon in Sec. 11.2.3. Thus the ternary Al-containing Laves phases in particular Nb(Ni,Al)2 or NbNiAl, Ta(Ni,Al)2 or TaNiAl, and Ta(Fe,Al)2 which have been mentioned in Sec. 11.8.1, have been considered for high-temperature applications because of their high melting temperature, high strength and potential oxidation resistance (Sauthoff, 1989, 1991a, 1992; Machon, 1992; Zeumer and Sauthoff, 1992; Von Keitz and Sauthoff, 1992). Again the monolithic Laves phases are too brittle for structural applications. However, these phases form stable, twophase equilibria with the B2 aluminides NiAl and FeAl, respectively, which offers

752

11 Intermetallics

the possibility of preparing intermetallic two-phase alloys with reduced brittleness. NiAl-NbNiAl, NiAl-TaNiAl, NiAl(Nb,Ta)NiAl, and FeAl-TaFeAl alloys have been prepared by ingot metallurgy methods and studied with respect to strength, ductility, toughness, brittle-to-ductile transition temperature, creep resistance, and oxidation resistance as has been discussed in Sec. 11.4.3.6 under the heading Structural Alloys.

11.9 Beryllides The transition-metal beryllides - in particular the Be-rich phases with Ti, Zr, Hf, Nb, Ta, or Mo - have long been regarded as very attractive for applications as structural, high-temperature materials because of their low densities between 2 and 5 g/cm3, high melting temperatures, high stiffnesses and strengths, and high oxidation resistances (Stonehouse et al., 1960; Ryba, 1967; Hove and Riley, 1967; Marder and Stonehouse, 1988; Tien et al., 1992; Kumar and Liu, 1993). Characteristic compositions of these Be-rich phases are MBe 13 for, e.g. M = Zr, or Hf, MBe 12 for, e.g. M = Ti, Nb, Ta, or Mo, and M 2 Be 17 for, e.g. M = Ti, Zr, Hf, Nb, or Ta. Such phases and other Be-rich phases with similar compositions are formed with most other metals. The multitude of intermetallic Be phases is similar to that of the intermetallic rare earth phases, which are addressed in the next section. No beryllides have been observed in the Al-Be and Al-Si systems (Massalski etal., 1990). These Be-rich phases have complex cubic, tetragonal, or hexagonal crystal structures, and it has to be noted that the crystallographic data in standard compilations are controversial in some cases (Massalski etal., 1990; Villars and Calvert, 1991).

Beryllides with less Be have simpler structures, e.g. the Laves phases MBe2 with M = Cr, Mo, or Fe with a C14 structure and M = Ti, Nb, or Ta with a C15 structure (see Sec. 11.8.1), the B2 phases NiBe and TiBe (see Sec. 11.4.4), or the A15 phase Mo3Be (see Sec. 11.7.1). Basic data can be found in early compilations (Shaffer, 1964; Samsonov and Vinitskii, 1980). The Be-rich phases with Nb, Zr, or similar transition metals have been prepared by powder metallurgy methods (Marder and Stonehouse, 1988; Henager etal., 1992 b), though ingot metallurgy methods have also been used (Nieh etal., 1990). Processing must be done with care because of the volatility and toxicity of Be (Stonehouse and Marder, 1990). The oxidation resistance is generally excellent due to the formation of protective BeO scales (Marder and Stonehouse, 1988; Grensing, 1989), though pest-like oxidation with specimen disintegration is observed in an intermediate temperature range aroung 800 °C for Nb and Zr beryllides (Westbrook and Wood, 1964; Ryba, 1967; Aitken, 1967; Chou et al., 1992). It is noted that beryllides have been proposed for applications as protective coatings, e.g. for Ta (Lawthers and Sama, 1993). The mechanical behavior of these Berich phases and its variation with temperature has been studied by means of hardness tests, bending stress-rupture tests, tension tests and compression tests (Ryba, 1967; Marder and Stonehouse, 1988; Fleischer and Zabala, 1990 c; Nieh and Wadsworth, 1990; Bruemmer etal., 1993). The observed brittle-to-ductile transition temperatures are of the order of 1000 °C. The low-temperature fracture toughness Klc has been found to be between 2 and 4 MN/m 3/2 with practically no macroscopic ductility (Bruemmer et al., 1993), though there are indications of local plasticity at

11.10 Rare-Earth Compounds

indentations (Ryba, 1967). The bending rupture strengths show a maximum above 1000 °C (Marder and Stonehouse, 1988; Henager et al., 1992b; Bruemmer etal., 1993) which may be due to pest-like oxidation and premature brittle fracture below 1000 °C. The elastic constants have been determined (Fleischer etal., 1989b) as well as the thermal expansion coefficients, thermal conductivities and specific heats (Marder and Stonehouse, 1988). Recent work has been concentrated on NbBe 12 with the tetragonal D2 b structure, and the deformation behavior has been studied in detail (Henager etal., 1992b). The dislocation reactions at high temperatures have been analyzed and the slip systems have been identified (Bruemmer etal., 1992b), the dislocation core structures have been studied by atomistic modeling (Shondi et al, 1992), and the twinning behavior has been analyzed (Chariot etal., 1991; Sondhi etal., 1993). The mechanical behavior of NbBe 12 may be improved by combining NbBe 12 with other metallic or intermetallic phases, e.g. Be, Nb 2 Be 17 , or NbBe 3 in the binary Nb-Be system (Bruemmer etal., 1990, 1992a). Likewise NbBe 12 , as well as TiBe 12 , ZrBe 13 , Nb 2 Be 17 , or Ta 2 Be 17 , have been proposed as the reinforcing phase in intermetallic matrix composites based on the B2 phase FeAl (Tien et al., 1992). In these and other composite systems chemical compatibility may be a problem since chemical reactions between the constituent phases affect the composite stability, as has been shown by combining NbBe 12 with other materials (Brimhall and Bruemmer, 1992) or TiBe12 with NiAl (Carbone etal., 1988). Besides these Be-rich phases, NiBe with the B2 structure is advantageous in view of its mechanical properties and its excellent oxidation resistance up to 1100°C (Lee

753

and Nieh, 1989; Nieh et al., 1990; Pharr et al., 1991). In particular it does not show pest-like oxidation at intermediate temperatures (Chou et al., 1992). It has an extended range of homogeneity (Massalski et al., 1990) which offers various possibilities for alloying. The simple B2 structure nourishes the expectation that the plastic deformation of NiBe is easier than that of the Be-rich phases. The room-temperature strength of NiBe shows a minimum at the stoichiometric composition, which is due not only to the strengthening effect of constitutional point defects for the off-stoichiometric compositions, as in the case of NiAl (see Sec. 11.4.3.3), but also to the strengthening effect of interstitially dissolved oxygen with a minimum solubility of O in NiBe at the stoichiometric composition (Nieh and Wadsworth, 1989; Nieh et al., 1989, 1990). Encapsulated NiBe can be hot forged and extruded at 1100°C (Nieh et al., 1990). In view of the attractive properties of NiBe and the Be-rich phases, much more work is necessary for the optimization of the mechanical behavior, i.e. for improving the low-temperature toughness and for overcoming the related processing problems.

11.10 Rare-Earth Compounds There is a multitude of intermetallic rare earth phases, since the rare earth metals have much larger sizes and lower electronegativities than most other metals, which gives many possibilities for space filling (Buschow, 1980). The complex crystal structures of these phases derive from a few basic structures, e.g. the cubic D2X structure or the hexagonal D2 d structure, by stacking the structural units in different ways from which the various stoichiometries result (Herget and Domazer,

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11 Intermetallics

1975; Buschow, 1980). The crystallographic data of most binary, ternary, and quaternary rare earth phases have been compiled (Gladyshevskii and Bodak, 1982). The rare earth intermetallics are mostly line compounds and have only very narrow composition ranges. Preparation of the phases is difficult because of the high affinity of the rare earths for oxygen, which leads to reactions with common crucible materials, and various processes have been developed for the preparation of the intermetallic rare earth phases (Herget and Domazer, 1975). 11.10.1 Magnet Materials The rare earth intermetallics are of great practical importance because of their outstanding magnetic properties, depending on the composition and crystal structure (Buschow, 1980, 1991). SmCo5 and Sm 2 Co 17 have given rise to the development of the rare earth permanent magnets (REPMs), which excel by their high energy product [B - H\max (B is the magnetic induction and H the magnetic field strength) and their high coercivity (Livingston, 1990; Buschow, 1991; Stadelmaier et al., 1991). Optimized property spectra have been achieved by alloying, i.e. substituting other rare earth metals for Sm and Fe and other transition metals as well as Cu for Co, and proper processing (Strnat, 1990). The resulting multinary materials have complex multiphase structures and the effects of the various structural features on the magnetic properties are not yet completely understood (Stadelmaier et al., 1988). Alloying with the interstitials B, C, and N has led to ternary rare earth phases with very attractive magnetic properties, and outstanding examples are Sm2Fe17CJC, Sm 2 Fe 17 N :c , Nd 2 Fe 1 4 B, and Nd 2 Fe 1 4 C

(Buschow, 1986, 1991). Nd 2 Fe 14 B has become the basis for the development of the N d - F e - B permanent magnets which have an even higher energy product [B • H\max at room temperature than the SmCo5-type and Sm2Co17-type materials at lower costs, whereas their behavior at elevated temperatures is less satisfactory (Strnat, 1990; Buschow, 1991). The various REPM materials, SmCo5-type, Sm 2 Co 17 type, and Nd 2 Fe 14 B-type, are produced primarily by powder metallurgy techniques though ingot metallurgy and rapid solidification have also been used (Herget and Domazer, 1975; Ormerod, 1988; Croker, 1990; Strnat, 1990; Buschow, 1991; Steinhorst, 1992). The materials are brittle, which makes machining difficult, and their corrosion resistance is lower, which necessitates protective coatings in corrosive environments. SmCo5 suffers an environmental degradation of coercivity which is due to hydrogen absorption (Willems and Buschow, 1987). Hydrogenation of the N d - F e - B alloys leads to structure transformation and decomposition reactions, and is used for producing fine-grained N d - F e - B powders (Harris and McGuiness, 1991). Because of their most advantageous combination of properties, the rare earth magnets have found a huge market with a broad field of applications (Strnat, 1990; Buschow, 1991). 11.10.2 Hydrogen Storage Materials Apart from applications as magnet materials, various rare earth phases can absorb large amounts of hydrogen and are of high interest for applications as hydrogen storage materials, as is exemplified by RCo 3 and RFe 3 (Christodoulou and Takeshita, 1993 a), or the rare earth Laves phases, which have been addressed in Sec. 11.8.2.3. The most outstanding exam-

11.11 Silicides

pie is given by the RNi 5 phases - in particular LaNi 5 which may be regarded as a prototype hydrogen storage material, and which is considered for applications as rechargeable electrodes in electrochemical nickel-hydrogen batteries (Reilly, 1979; Willems and Buschow, 1987). The crystal structure of the ternary hydride LaNi 5 H x and the similar RCo 5 H x is complex and various models have been proposed for the distribution of the H atoms which are still in discussion (Somenkov and Shilstein, 1979; Willems and Buschow, 1987). The stability of LaNi 5 is affected by absorption-desorption cycling which is due to surface degradation processes (see e.g. Uchida etal., 1991; Josephy etal., 1991). The property spectrum has been optimized by alloying with further elements as is exemplified by the development of La 0 8 Nd 0 2 Ni 2 5 Co 2 4 Si 0 ! (Willems and Buschow, 1987) or MmNi 3 5 Co 0 7A1O 8 where Mm stands for mischmetall, which is a mixture of rare earth metals (Sakai et al., 1991). The permanent-magnet phase SmCo5 can also absorb large quantities of hydrogen which, however, reduces its coercive force (Buschow, 1991). It is noted that CaNi 5 , which is isostructural with LaNi 5 , shows an analogous behavior and is considered as a low-cost candidate phase for hydrogen storage (Yagisawa and Yoshikawa, 1979).

11.11 Silicides The transition metal silicides show close similarities to the intermetallics and thus they are frequently classed with the intermetallics though silicon is no longer a metal, but a semiconductor. Silicides are generally hard and brittle with a metallic luster, high electrical and thermal conductivities, a positive temperature coefficient of

755

resistivity, and paramagnetism (Westbrook, 1960 b; Nowotny, 1963; Wehrmann, 1967; Goldschmidt, 1967 b). In detail the type of bonding and the electrical conductance depend on the metal-to-silicon ratio and on the particular metal M. The highest metallicity is manifested by the metal-rich silicides M3Si and M2Si, in which the Si atoms are isolated. The lower the M/Si ratio, the lower the metallicity, and the more readily continuous Si chains or networks are formed in the crystal lattice (Goldschmidt, 1967 b). However, it has to be emphasized that the situation is rather complex, e.g. the disilicide CrSi2 with the hexagonal C40 structure is indeed a semiconductor, whereas NbSi 2 with the same crystal structure and MoSi2 with the C40 structure at very high temperatures and the closely related tetragonal C l l b structure at lower temperatures (Massalski etal., 1990) are metallic (Nowotny, 1963, 1972 a; Goldschmidt, 1967 b). Likewise FeSi and CoSi - both with the cubic B20 structure - form a continuous series of solid solutions with a transition from semiconduction for the Fe-rich phases to metallic conduction for the Co-rich phases (Romasheva et al., 1980), which affects the magnetic properties and the elastic behavior (Zinov'eva et al., 1984). Silicides were considered for high-temperature applications 40 years ago because of their high strength and oxidation resistance at high temperatures (Lowrie, 1952; Fitzer, 1952; Schwarzkopf and Kieffer, 1953; Arbiter, 1953 b; Wehrmann, 1956). This interest initiated extended research activities with a screening of the physical and mechanical properties, and the results have been the subject of detailed reviews and data compilations (Westbrook, 1960 b; Shaffer, 1964; Hove and Riley, 1967; Wehrmann, 1967; Goldschmidt, 1967 b; Samsonov and Vinitskii, 1980). The phase

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11 Intermetallics

diagrams of ternary systems of interest have been analyzed (e.g. Kudielka and Nowotny, 1956; Brukl etal., 1961; Schob etal., 1962; Bardos etal., 1966; Gladyshevskii and Borusevich, 1966; Skolozdra et al., 1966 a, b), as well as the crystal structures and the crystal chemistry of the various silicides (Nevitt, 1963,1967; Nowotny, 1963, 1972 a; Nowotny and Benesovsky, 1967; Jeitschko etal., 1969; Jeitschko, 1977). Only one successful development emerged from these early efforts, and it led to the application of MoSi 2 as heating elements in high-temperature furnaces (Fitzer and Rubisch, 1958; Tamura, 1961; Schrewelius and Magnusson, 1966; Schlichting, 1986). Later, protective coatings on metallic alloys made use of the advantageous oxidation resistances of various silicides (e.g. Hildebrandt et al., 1978; Fitzer et al., 1978; Caillet etal., 1978; Fitzer and Schlichting, 1983; Meier, 1987; Packer, 1989). The advantageous electrical properties of various silicides have led to important applications as thin films in microelectronic devices (Nicolet and Lau, 1983; Murarka, 1983 a, 1984). In view of such applications, the electrical resistivities, thermodynamical properties, thin-film formation kinetics and diffusivities of the transition-metal silicides were reviewed (Murarka, 1983 b). Presently, the need for new structural materials for the highest temperatures has revived the interest in silicides for structural applications, and respective developments are under way (Kumar and Liu, 1993). In the following sections a brief overview is given on the important features of those silicides which have already been applied or are being considered for applications.

11.11.1 M3Si Phases Cr3Si, which is a candidate phase for high-temperature applications, and V3Si, which is superconducting at low temperatures with a comparatively high transition temperature, both crystallize with the topologically close-packed cubic, A15 structure and have been discussed in Sees. 11.7.7 and 11.7.2, respectively. A metal-rich silicide with a very simple crystal structure is Ni3Si with the ordered f.c.c. LI 2 structure. It has a very high potential for structural applications because of its advantageous mechanical properties and its outstanding corrosion resistance. Its deformation behavior is similar to that of other Ll 2 phases, in particular Ni 3 Al, and thus Ni3Si has been discussed in Sec. 11.4.2.2 together with other L l 2 phases. The ordered b.c.c. D0 3 structure is adopted by Fe3Si which shows close similarities to Fe3Al (see Sec. 11.5.1). According to the generally accepted phase diagram (Schlatte and Pitsch, 1976; Kubaschewski, 1982), the D0 3 structure of Fe3Si is stable up to about 800 °C (for the stoichiometric composition) where it transforms to the B2 structure, and at about 1000 °C Fe3Si disorders to form the b.c.c. solid solution. The existence of the B2 structure has been questioned on the basis of neutron scattering results (Hilfrich etal., 1990) which, however, are not unambiguous (Inden, 1993). The deformation behavior of Fe3Si has been analyzed in detail to study the effects of atomic order on the mechanical properties (Saburi etal., 1968; Lakso and Marcinkowski, 1974; Marcinkowski, 1974 a; Ehlers and Mendiratta, 1984; Oertel etal., 1986). Elastic moduli and dislocation line energies have been determined (Kotter et al., 1989) as well as vacancy formation ener-

11.11 Silicides

gies (Schaefer etal., 1992), and diffusion mechanisms (Sepiol and Vogl, 1993). The electron distribution in Fe3Si has been studied with respect to the effects of ordering on the magnetic behavior (Himsel et al., 1980; Blau et al, 1980). The ternary Fe3(Si,Al) exhibits a high magnetic permeability and has found widespread application as a magnetic recording-head material known as Sendust alloy (Yamamoto, 1980; Brock, 1986). The magnetic properties can be optimized by proper alloying (Yamamoto and Utsushikawa, 1977; Miyazaki etal., 1992), and the processing problems due to brittleness have been overcome (Watanabe et al. 1984; Shao et al., 1991). Fe3Si has also been applied as a structural material because of its excellent corrosion resistance (Liu etal., 1990; Lou etal., 1991; Kumar and Liu, 1993). 11.11.2 M2Si Phases M2Si phases may be present in protective coatings (Hildebrandt et al., 1978; Meier, 1987). High-temperature structural applications are expected for the lightweight phase Mg2Si which is regarded as highly promising for use as pistons in car engines (Schmid etal., 1990; von Oldenburg etal., 1990). Mg2Si crystallizes with the f.c.c. Cl structure with 12 atoms per unit cell, and its density is only 1.88 g/cm2. The mechanisms of slip, twinning, and cleavage have been discussed (Paufler and Schulze, 1978). Mg2Si has a comparatively high strength and a low thermal expansion coefficient. However, its brittleness with a brittle-to-ductile transition temperature of 450 °C precludes the use of single-phase Mg2Si. The methods of induction melting have been studied in detail for various Mg2Si alloys (G. H. Li etal., 1993). The ternary system Mg-Si-Al (Liidecke, 1986) offers various possibilities for alloy-

757

ing in order to combine Mg2Si with a second ductile phase which may relieve the brittleness. Indeed alloying with Al produces Al-Mg 2 Si alloys with brittle Mg2Si particles in a ductile Al matrix, which has given rise to a successful on-going alloy development (Schmid etal., 1990; von Oldenburg et al., 1990). The effect of the distribution of particles on the deformation behavior has been studied for model Al-Mg 2 Si alloys (Liu, 1989). Mg2Si is a semiconductor with a high thermoelectric power and a low thermal conductivity which is advantageous for thermoelectric power generation (Noda et al., 1992 a, b). The thermal conductivity can be minimized by alloying with Ge to form the ternary Mg2(Si,Ge) phase. Doping with Ag and Sb produces p-type and n-type semiconducting Mg2(Si,Ge), respectively, and the combination of both results in a thermocouple with a high efficiency for thermoelectric energy conversion. The interactions between doping, point defects and microstructure have been studied in detail (Andreyeva et al., 1988). Metal-like M2Si phases, as well as other transition metal silicides, are of great importance for applications as thin-film materials in electronic devices, i.e. as low-resistivity interconnects and gates, ohmic contacts, and Schottky barriers in very-large-scale-integrated circuits (VLSI) (Nicolet and Lau, 1983; Murarka, 1983 a, 1984). The most eminent example is Pd2Si which is used for shallow contacts with very small, controlled thickness and extension (Chapman et al., 1979; Kritzinger and Tu, 1981). The kinetics of thin-film formation and growth are controlled by diffusion and can be optimized by alloying with further elements (Mayer et al., 1979; Olowolafe etal., 1979; Tu etal., 1980; Eizenberg etal., 1981; Eizenberg and Tu,

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11 Intermetallics

1982). The film formation process affects the distribution of dopants in the Si substrate (Wittmer and Seidel, 1978; Kikuchi, 1983; Wittmer et al., 1983). The rate controlling mechanisms of the film formation process with nucleation and growth have been analyzed in detail (d'Heurle and Gas, 1986; d'Heurle, 1993 a, b). Other M2Si phases have also been studied with respect to thin-film formation, e.g. Ni2Si (d'Heurle etal., 1984; Aly and Stark, 1984), Co2Si (Hattori etal., 1988; Chen and Chang, 1993), or Mg2Si (Lim and Stark, 1984). Diffusion in thin films has been analyzed and compared with bulk diffusion (Tu et al., 1983). Thin oxide layers may form on the silicide films by air exposure at room temperature which can be a problem for ohmic contacts or a necessity for a tunneling device (Cros, 1983). The partial substitution of the transition element M in M2Si by a second transition metal M' leads to ternary silicides of approximate composition MM'Si corresponding to (M,M')2Si. Such ternaries are primarily the Si-containing E phases and V phases (Jeitschko et al., 1969; Jeitschko, 1970) and the ternary Si-containing Laves phases (Bardos et al., 1961), which were discussed in Sec. 11.8, as well as many other phases, which all differ by composition and crystal structure (Nowotny, 1972 a). This is exemplified by the Fe-Nb-Si system with the ternary silicides E, V, x1? x2, x3 and the Laves phase Nb(Fe,Si)2 with up to 25 at.% Si (Raghavan, 1987), or the Co-Nb-Si system with the ternary silicides E, T, v, r|, \|/, and the ternary Laves phase Nb(Co,Si)2 with Si contents between about 10 and 20 at.% (Argent, 1984). Finally, it is noted that other phases - in particular a phases and A13 Mn-base phases - dissolve large amounts of Si by which these phases are stabilized (Gupta etal., 1960; Bardos etal., 1966).

11.11.3 M 5 Si 3 Phases The M 5 Si 3 phases adopt a variety of complex crystal structures (Franceschi and Ricaldone, 1984). The transition-metal M 5 Si 3 phases crystallize primarily with the tetragonal D8 m structure, e.g. for M = V, Cr, Mo, or W, or the hexagonal D8 8 structure, e.g. for M = Mn, or Ti (Franceschi and Ricaldone, 1984; Massalski etal., 1990). Most of these phases are very stable, congruently melting line compounds with melting temperatures above 2000 °C. However, Fe 5 Si 3 with a D8 8 structure is stable only between 825 and 1060 °C, and Cr 5 Si 3 with a D8 m structure undergoes a polymorphic transformation above 1500°C (Massalski et al., 1990). Nb 5 Si 3 and Ta5Si3 exhibit the D8 m structure at low temperatures and the likewise tetragonal D8t structure at high temperatures (Parthe etal., 1955; Massalski etal., 1990). Various tetragonal M 5 Si 3 phases, e.g. for M = V, Nb, Ta, Cr, Mo, or W, as well as Hf5Si3, crystallize with the D8 8 structure if they contain small amounts of interstitial impurities, in particular carbon (Parthe etal., 1955; Goldschmidt, 1967b; Franceschi and Ricaldone, 1984). Such a stabilization effect of interstitial impurities has also been observed for other intermetallics (Gschneidner, 1993). The most attractive M 5 Si 3 phase for structural high-temperature applications is Ti5Si3 with the hexagonal D8 8 structure because of its high stability with a melting temperature of 2130°C, its high strength and hardness, and its low density (Liu etal., 1988; Beaven etal., 1989; Rosenkranz et al., 1992; Kumar and Liu, 1993). Density data between 4.0 and 4.5 g/cm3 have been reported (Shaffer, 1964; Beaven etal., 1989; Fleischer and Zabala, 1990a; Frommeyer et al., 1990) as well as a Young's modulus of about 150 GPa (Beav-

11.11 Silicides

en etal., 1989; Frommeyer etal., 1990) and an anisotropic thermal expansion coefficient of about 6 x 10~ 6 /K (Frommeyer etal., 1990; Nakashima and Umakoshi, 1992). The outstanding strength of Ti 5 Si 3 is correlated with a very low fracture toughness and a brittle-to-ductile transition temperature above 1000 °C (Liu et al., 1988; Frommeyer etal., 1990; Vehoff, 1992; Vehoff etal., 1993). The oxidation resistance of Ti5Si3 is not as high as that of NiAl, but higher than that of TiAl (Liu etal., 1988; Murata etal., 1991; Thorn etal., 1993). As with other intermetallics, it is expected that the brittleness of Ti 5 Si 3 can be relieved by combining it with other less brittle phases. The crack tendency in Ti 5 Si 3 can indeed be reduced by reducing the Si content and alloying with Cr and Zr to produce three-phase, Ti5Si3-base alloys (Liuet al., 1988). Eutectic, unidirectionally solidified Ti-Ti 5 Si 3 alloys show advantageous combinations of strength, creep resistance, and fracture toughness (Crossman and Yue, 1971; Frommeyer etal., 1990). A promising materials development is based on the combination of Ti3Al and Ti5Si3 (Wu etal., 1989). Such Ti 3 AlTi5Si3 alloys, which are further alloyed with Nb, and similar alloys have been studied with respect to phase equilibria (X S. Wu etal., 1991), microstructure-property relationships (Wagner etal., 1991), creep (Es-Souni et al., 1992a, d) and fracture toughness (Vehoff, 1992, 1993). Besides ingot metallurgy, powder metallurgy methods with gas-atomization (Es-Souni et al., 1992b, c) have been used as well as mechanical alloying (Calka etal., 1991), and combustion synthesis has also been considered (Bhaduri, 1992). Ti 5 Si 3 has been proposed as a constituent phase in composites (Kumar, 1991; Shah and Anton, 1992 b), i.e. as the matrix phase with

759

ceramic reinforcements (Meschter and Schwartz, 1989; Bhattacharya, 1991) or as the reinforcing phase in an MoSi2 matrix (Wiedemeier and Singh, 1992; Schwartz etal., 1993). Another candidate phase for high-temperature applications is Nb 5 Si 3 with a melting temperature of 2484 °C (Massalski et al., 1990), which is still more stable than Ti 5 Si 3 and which shows the above-mentioned polymorphism. Again its brittleness precludes the application of single-phase Nb 5 Si 3 . However, the reduction in the Si content leads to two-phase Nb 5 Si 3 -Nb alloys with ductile Nb-rich particles in the brittle Nb 5 Si 3 matrix (Lewandowski et al., 1988). These alloys show improved toughness because of crack-bridging (Rigney etal., 1991; Mendiratta etal., 1991; Mendiratta and Dimiduk, 1993). The phase relationships in these alloys and the effects of rapid solidification have been studied (Cockeram etal., 1991; Bertero etal., 1991b). Nb 5 Si 3 is subject to pest-like oxidation (Westbrook and Wood, 1964), and Nb 5 Si 3 -Nb alloys are embrittled by oxygen and hydrogen exposure (Rigney et al., 1992). Other M 5 Si 3 phases are less attractive for applications, and thus they have been studied less. M 5 Si 3 phases can be synthesized by mechanical alloying, which leads to amorphization in the case of Ta5Si3 (Kumar and Mannan, 1989). Y 5 Si 3 and the ternary Y5(Si,Ge)3 are of interest for use as hydrogen storage materials (McColm and Ward, 1992). 11.11.4 MSi Phases The transition metal silicides CrSi, MnSi, FeSi, and CoSi crystallize with the cubic B20 structure whereas NiSi and the noble metal silicides PtSi, IrSi, and PdSi crystallize with the orthorhombic B31

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structure (Massalski etal., 1990). MSi phases - in particular PtSi and also NiSi are important thin-film materials for applications in electronic devices (see, e.g. Eizenberg etal., 1981; Cohen etal., 1982; Murarka, 1984; Tu etal., 1983; Appelbaum et al., 1984). These phases have been studied in close contact with the work on Pd2Si, which was addressed in Sec. 11.11.2, and the respective cited publications are usually directed at the ensemble of thinfilm silicides of interest. Besides electronic applications, NiSi is of interest as a constituent phase in oxidation resistant coatings (Meier, 1987) and as a heat storage material with high thermal conductivity (Wilson and Cavin, 1992). IrSi forms in MoSi2/Ir layers which are considered for the coating of carbon-carbon composites (Chou, 1990; Chou and Nieh, 1990). MnSi and CoSi exhibit large thermoelectric powers (Samsonov and Vinitskii, 1980) and thus are promising for thermoelectric power generation (Sakata and Nishida, 1976). In particular, a CoSi-CrSi 2 thermogenerator has been proposed (Sakata and Tokushima, 1963). 11.11.5 Disilicides The transition metal MSi2 phases with M = Ti, Zr, Hf, V, Nb, Ta, Cr, Mo, or W crystallize with the tetragonal C l l b structure, the hexagonal C40 structure or the orthorhombic C49 and C54 structures and show extended mutual solid-solubilities (Kudielka and Nowotny, 1956). These crystal structures are closely related since they only differ in the different stacking of the same structural elements (Nowotny, 1963). MoSi2 with a C l l b structure has found widespread use as heating elements in high-temperature furnaces for temperatures up to 1700°C because of its advanta-

geous electrical properties and its excellent oxidation resistance (Fitzer and Rubisch, 1958; Tamura, 1961; Schrewelius and Magnusson, 1966; Schlichting, 1986). This high oxidation resistance, which is not shown by the other Mo-rich silicides, is due to the formation of a protective, vitreous, highly adherent SiO2 film with prior evaporation of the volatile Mo oxide (Fitzer, 1955; Kieffer and Benesovsky, 1956; Lee etal., 1991; Meier and Pettit, 1992). However, MoSi2 is subject to oxidative disintegration by pest-like oxidation at intermediate temperatures between about 300 °C and 600 °C (Fitzer, 1955; Westbrook and Wood, 1964; Chou and Nieh, 1993 a), which occurs at grain boundaries and micro structural defects and can thus be minimized by avoiding grain boundaries and microstructural defects (Chou and Nieh, 1993 b) or by alloying with small amounts of Fe and Re (Ban and Ogilvie, 1966). This high oxidation resistance makes MoSi2 attractive for use as a protective coating (Fitzer, 1952; Motojima et al., 1982; Packer, 1989; Chou, 1990; Petrovic, 1993). Apart from these high-temperature applications, MoSi2 is advantageous for applications in electronic devices - either as the substrate for Si films (Campisi et al., 1981) or as thin-film interconnectors in VLSI circuits (see e.g. Chow et al., 1980; Murarka, 1984; Urwank etal., 1985; and compare with Sec. 11.11.2). The kinetics of silicide formation in thin films have been analyzed in detail (d'Heurle, 1993 b). Presently MoSi2 is being considered for structural, high-temperature applications and various materials developments are in progress (Lugscheider et al., 1991; Kumar and Liu, 1993; Petrovic, 1993; Hardwick etal., 1993). The elastic constants are known (Nakamura et al., 1990; Nakamura, 1991; Srinivasan and Schwarz, 1992), and the plastic deformation behavior

11.11 Silicides

has been analyzed with respect to dislocation configurations and slip systems (Umakoshi etal., 1990a, b; Maloy e t a l , 1992; Evans et al, 1993; Rao et al., 1993), twinning (Mitchell et al., 1992), and creep (Tamura, 1961; Kimura etal., 1990; Sadananda and Feng, 1993). The brittleto-ductile transition occurs at about 1000 °C or higher temperatures depending on the microstructure and the impurity content (Umakoshi etal., 1991; Aikin, 1992; Srinivasan etal., 1993; Petrovic, 1993). The ductility and toughness of MoSi2 can be improved by special processing routes (e.g. Tiwari et al., 1991; Castro etal., 1992; Patankar and Lewandowski, 1993), microalloying with C to avoid SiO2 layers on the grain boundaries (Maloy etal., 1991; Jayashankar and Kaufman, 1992), the application of soft, dislocation emitting surface films (Czarnik et al., 1993), or the insertion of a ductile second phase which is usually a Nb-rich phase (Lu etal., 1991; Xiao and Abbaschian, 1992; Venkkateswara Rao etal., 1992; Alman and Stoloff, 1992). Toughness can also be increased by embedding hard second phases, in particular ceramics, in an MoSi2 matrix in order to hinder crack growth (see e.g. Bhattacharya and Petrovic, 1991; H. Chang etal., 1992). In view of the beneficial effects of second phases, combinations of MoSi2 with other phases are being studied with the aim of developing new MoSi2-based composites (Meschter and Schwartz, 1989; Shah et al., 1990; Yang et al., 1990; Boettinger et al., 1992; Petrovic and Vasudevan, 1992; Wiedemeier and Singh, 1992; Petrovic, 1993; Sadananda and Feng, 1993). Most work has been concentrated on the MoSi 2 -SiC system (Bhattacharya and Petrovic, 1991; Henager et al., 1992a; Jayashankar and Kaufman, 1992; Suzuki et al., 1992; Alman and Stoloff, 1993; Feng

761

and Michel, 1993; Jeng et al., 1993; Ting, 1993). Other systems of particular interest are MoSi 2 -TiC (Yang and Jeng, 1990; Chang and Gibala, 1993) and MoSi 2 A12O3 (Alman et al., 1991; Y S. Kim et al., 1991; Meschter, 1991), as well as the eutectics MoSi 2 -Mo 5 Si 3 (Mason and Van Aken, 1993) and MoSi 2 -Er 2 Mo 3 Si 4 (Patrick and Van Aken, 1993) and the ductile-phase toughened MoSi 2 -Nb/Ta (Carter and Martin, 1990; Lu et al., 1991; Alman and Stoloff, 1992; Venkkateswara Rao et al., 1992), and MoSi 2 -Mo systems (Deve etal., 1992). Furthermore, MoSi2 has been considered as a reinforcing phase in other composite systems, e.g. in an SiC matrix (Lim et al., 1989) or in a beryllide matrix (Bruemmer et al., 1993). WSi2 with a tetragonal C l l b structure is another disilicide with a high oxidation resistance which relies on the formation of a vitreous, dense, adherent SiO2 scale at high temperatures, as in the case of MoSi2 (Kieffer and Benesovsky, 1956). At intermediate temperatures pest-like oxidation occurs (Westbrook and Wood, 1964). The elastic constants of WSi2 (Nakamura, 1991) and the plastic deformation at high temperatures (Kimura et al., 1990) have been studied. Apart from possible hightemperature applications (Kumar and Liu, 1993), WSi2 is of interest for thin-film applications in electronic devices (see, e.g. Olowolafe etal., 1979; Murarka. 1984). TiSi2 with an orthorhombic C54 structure is protected against oxidation by a vitreous, adherent TiO 2 -SiO 2 scale (Kieffer and Benesovsky, 1956; Rahmel and Spencer, 1991). In view of its high oxidation resistance and its low density, TiSi2 has been considered for structural applications at high temperatures (Lugscheider etal., 1991; Rosenkranz etal., 1992; Kumar and Liu, 1993). It can be synthesized by mechanical alloying (Calka et al.,

762

11 Intermetallics

1991; Matteazzi et al., 1992), single crystals have been prepared (Thomas et al., 1987; Peshev etal., 1989), and the mierostructure has been analyzed (Jia et al., 1989). Data for thermal expansion, elastic moduli, hardness, yield strength, creep resistance, and toughness have been obtained (Lugscheider etal., 1991; Rosenkranz and Frommeyer, 1992; Vehoff, 1993). Apart from eventual high-temperature applications (Kumar and Liu, 1993), TiSi2 is of interest for thin-film applications in electronic devices (see, e.g. Murarka, 1984). ZrSi2 with the orthorhombic C49 structure shows attractive physical and mechanical properties with respect to structural high-temperature applications (Rosenkranz and Frommeyer, 1992). Like other disilicides, it has a low fracture toughness with a brittle-to-ductile transition temperature of about 900 °C (Vehoff, 1993). A ZrSi2 coating on Zr provides good oxidation resistance (Caillet et al., 1978) though bulk ZrSi2 does not form dense, adherent scales on oxidation (Kieffer and Benesovsky, 1956). The disilicides VSi2, NbSi 2 , and TaSi2 crystallize with the hexagonal C40 structure (Kudielka and Nowotny, 1956). The elastic constants of VSi2 have been determined (Fleischer et al., 1989 b; Nakamura, 1991), as well as its hardness and apparent brittle-to-ductile transition temperature (Fleischer etal., 1990). VSi2 is of interest for applications in electronic devices and its thin-film formation reactions have been studied (Lim and Stark, 1984). TaSi2 has been successfully used in producing microprocessors and other electronic devices (Murarka, 1984). The mechanisms of thinfilm formation of TaSi2 have been studied in detail (see, e.g. Maa et al., 1985; Natan, 1985; Nava et al., 1985). CrSi2 with a hexagonal C40 structure is being considered for high-temperature ap-

plications (Kumar and Liu, 1993) though it does not form dense, adherent, protective oxide scales at 1200°C (Kieffer and Benesovsky, 1956; Grabke and Brumm, 1989). Single crystals have been prepared (Peshev et al., 1989), and the elastic properties (Nakamura, 1991), and the plastic deformation behavior have been studied (Umakoshi et al., 1991). CrSi2 is of interest for applications as thin-film interconnectors in electronic devices (Tu et al., 1980; Appelbaum et al., 1984). It exhibits a large thermoelectric power (Samsonov and Vinitskii, 1980; Nishida and Sakata, 1978), and it has been proposed for application as thermoelectric material, e.g. in a CrSi 2 CoSi thermogenerator for thermoelectric power generation (Sakata and Tokushima, 1963; Ohkoshi et al., 1988). FeSi2 is another disilicide with a high thermoelectric power and has been proposed for use in a thermoelectric power generator (Hesse, 1969 b). It has an orthorhombic crystal structure and is stable only at lower temperatures, whereas at higher temperatures it decomposes to form a two-phase FeSi-Fe 2 Si 5 alloy (Nishida, 1973; Kojima et al., 1990). Furthermore, it is subject to pest-like oxidation (Westbrook and Wood, 1964). Doping with Al produces p-type conducting FeSi2, doping with Co produces n-type FeSi2, and the combination of both materials gives a thermocouple which can be used as a thermoelectric power generator (Hesse, 1969 a). CoSi 2 , which crystallizes with the comparatively simple cubic Cl structure, is applied in electronic, large scale integration devices and the conditions for thin-film formation have been studied (see, e.g. Pretorius etal., 1985; Catana etal., 1992; Mantl, 1993; Kumar, 1994). In view of its simple crystal structure, CoSi2 is expected to show plastic deformability and is considered for structural applications (Ya-

11.12 Prospects

maguchi et al., 1993). The Co-Si phase diagram has been analyzed (Ishida et al., 1991b; Choi, 1992), the electron distribution has been studied by ab initio calculations (Sen Gupta and Chatterjee, 1986), the formation enthalpy of vacancies has been determined by positron annihilation (Ito et al., 1993), an excellent oxidation resistance has been found (Anton and Shah, 1989), and the plastic deformation behavior has been analyzed in detail with respect to dislocation configurations and slip systems (Takeuchi et al., 1991,1992; Ito et al., 1992; Suzuki and Takeuchi, 1993; Yamaguchi etal., 1993; Anongba and Steinemann, 1993). NiSi2 is of primary interest for applications in electronic thin-film devices (Kumar, 1994). In view of these applications, the conditions for thin-film formation have been studied in detail for NiSi2 (see, e.g. Tu etal., 1983; d'Heurle etal., 1984; Singh and Khokle, 1987) and likewise the formation of thin oxide layers on NiSi2 (Cros, 1983). Apart from these applications, the termal expansion behavior of NiSi2 has been studied with respect to use as a heat storage material (Wilson and Cavin, 1992). Finally, it is noted that the alloying of NiSi2 with Ni and Al produces phases with distinct color - pale blue, yellow, white, and variations of these - which may be of interest for applications in jewellery (Cortie et al., 1991).

11.12 Prospects As already stated in the introduction to this chapter, only a restricted selection of intermetallics could be assessed here. An enormous multitude of intermetallic phases is known (Villars and Calvert, 1991), and there are many phases which were not mentioned in the preceding sections and

763

which are of interest, e.g. as strengthening or embrittling phases in commercial metallic alloys. This is illustrated by the topologically close-packed |i, a, %, and related phases (Nevitt, 1963; Bardos etal., 1966; Benjamin etal., 1966; Hall and Algie, 1966; Sinha, 1973) which pose problems in many metallic alloy systems because of their embrittling effects, but which may also be used as strengthening phases if adequately distributed (Pickering, 1976; Bhandarkar etal., 1976; Gaspard etal., 1977; Sauthoff and Speller, 1982; Schumacher and Sauthoff, 1987; Sha etal., 1993). Most of the less known phases have not been studied sufficiently to establish a database which allows estimates of their potential for applications. Thus there is a great need for the determination of the basic properties - in particular constitution, phase diagrams, densities, electrical and thermal conductivities, elastic constants, and diffusion coefficients - of the less known intermetallics. As to the better known phases, which have been selected for materials developments or are regarded as candidate phases for new developments, the physical understanding of their properties and their behavior during processing and service is still insufficient, which impedes the progress of the present developments. Structural alloy developments usually result in complex multiphase intermetallic alloys which are optimized with respect to phase composition and phase distribution in order to obtain high strengths and corrosion resistances with sufficient formability. However, the effects of alloying elements on the basic properties of the respective phases are not yet understood sufficiently for these intermetallic alloys. One needs to know in what way a composition change in a particular phase changes, e.g., the elastic moduli. It is hoped that more experimental

11 Intermetallics

just the property spectra to the specific applications. The much advanced nickel aluminides and titanium aluminides can be used only up to about 1000°C because of their limited strength or oxidation resistance or both at higher temperatures, as has been stated before (Sauthoff, 1994). For applications significantly above 1000 °C other less-common phases with higher melting temperatures have to be used. Such phases are available, and examples are shown in Fig. 11-34 (Sauthoff, 1992). In comparison to the nickel aluminides and titanium alu-

Ti5S.3\HPSN \ \

25 c

i \ i \

NbNiAI-NiAl

E

NK

20-

res

work will be done with respect to the alloy systems of interest to obtain the missing data, and that more theoretical work - including quantum-mechanical, ab initio calculations - will be done to understand why the properties change with specific alloying additions in specific ways. This lacking understanding is not a characteristic of intermetallics alone. However, the problems of understanding are more pronounced for the intermetallics since the properties of the intermetallics vary with changing composition - as does the bonding character to a larger degree than is known for the familiar metallic alloy systems. As to applications, various functional intermetallic alloys are well established and are successfully used, e.g. magnetic FeCo alloys or superconducting A15 materials. The limited or lacking deformability poses problems for the processing of these materials, which is the subject of continuous research and development. A lot of progress has been achieved by applying advanced technologies which have been developed for other groups of materials. In particular, experiences in the processing of brittle structural intermetallics being developed now, and of the still more brittle advanced ceramics, may be useful here. New structural intermetallic alloys for high-temperature applications are at the center of the present interest in intermetallics, which is still growing. A few developments, which are based on the classic phases Ni 3 Al, Ti3Al and TiAl, and which are known as the nickel aluminides and the titanium aluminides, are on the brink of commercialization, but even these developments are still at an early stage compared with other developments of advanced materials, e.g. the modern engineering ceramics. Much more experimental and theoretical work is necessary to solve the processing problems and to ad-

MA 6000

^ \

15

AI3Nb\

ieli

764

\ • \ TaFeAl

4

\ ) 2 T i A I \ \ \\\

u 4—

10-

I to O

5 \\

°(

\

A

)

500

1000

1500

temperature in °C Figure 11-34. Specific yield strength (0.2% proof stress in compression per unit weight density at 10 ~ 4 s~1 strain rate) as a function of temperature for the D0 22 phase Al3Nb (Reip, 1991; Reip and Sauthoff, 1993), the Heusler-type phase Co2TiAl (Sauthoff, 1990b), the Laves phases TiC^ 5 Si 0 5 and TaFeAl (Sauthoff, 1990 b; Machon, 1992), the two-phase alloy NbNiAI-NiAl with 15 vol.% NiAl in the Laves phase NbNiAl (Sauthoff, 1990 b; Machon, 1992), and the hexagonal D8 8 phase Ti5Si3 (Frommeyer et al., 1990) compared with the superalloy MA 6000 (in tension) (Inco, 1982) and hot-pressed silicon nitride HPSN (upper limit of flexural strength) (Porz and Grathwohl, 1984).

11.13 Acknowledgements

minides, the less-common phases are stronger and more brittle, their crystal structures are complex, their handling is difficult, and thus they are regarded as exotic. However, these exotic phases may fill the gap between metallic high-temperature alloys and ceramics, as is clear from Fig. 11-34. The brittleness of the less-common phases can be alleviated by combining them with softer phases to form multiphase alloys with adequate microstructures. Even strengthening hard phases may improve the toughness by impeding crack growth. The mechanical behavior can be optimized by optimizing the microstructure, which requires careful control of the processing. However, it has to be emphasized that one cannot expect to obtain new intermetallic materials with properties similar to existing conventional metallic alloys. The "ductilizations" that have been achieved in a few cases - in particular Ni3Al and (Fe,Co,Ni)3V - rely on rather specific mechanisms and cannot be expected for other intermetallic phases. Thus intermetallic materials have to be regarded as a materials class of their own with property spectra which differ significantly from those of other materials and which can be varied within broad limits corresponding to metals on one side and nonmetals on the other side. This offers enormous possibilities for manifold developments which are exciting with respect to both practical applications and materials science. Finally, it is noted that much interest is concentrated on the development of intermetallic alloys for use as blades in flying gas turbines. This application is most demanding and it is not clear whether all the problems with strength, ductility, toughness, and corrosion resistance can be solved at economic costs. Less-high-technology applications may be more reward-

765

ing at present for the introduction of new intermetallic materials. An example may be the car engine, where strong, light components with sufficient corrosion resistance are needed. Here brittleness is not the problem since designers have learnt to use ceramic materials, e.g. for valves. However, new materials must be compatible with the metallic engine with respect to the physical properties, in particular thermal expansion and thermal conductivity. This requirement corresponds to the characteristics of intermetallic materials which are hard and brittle with mainly metallic atomic bonding, i.e. metallic, physical properties. Thus new intermetallic materials are expected to play an important role in the manufacture of car engines and similar applications.

11.13 Acknowledgements The author's research on intermetallics has been supported by the Deutsche Forschungsgemeinschaft (DFG) and the Bundesminister fur Forschung und Technologie (BMFT) throughout a decade, and this is gratefully acknowledged. The author is indebted to numerous colleagues whose researches are acknowledged in the references - in particular to his colleagues at the Max-Planck-Institut fur Eisenforschung - for many valuable and stimulating discussions. The author thanks Mrs Erika Bartsch for a great deal of electron microscope work on various intermetallic alloys and Mr Gerhard Bialkowski for innumerable mechanical tests on rather brittle materials.

766

11 Intermetallics

11.14 References

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Yang, W. S., Mikkola, D. E. (1993), Scr. Metall Mater. 28, 161-165. Yasuda, H., Takasugi, X, Koiwa, M. (1991a), in: Proc. Int Symp. on Intermetallic Compounds Structure and Mechanical Properties (JIMIS-6): Izumi, O. (Ed.). Sendai: The Japan Institute of Metals, pp. 33-37. Yasuda, H., Takasugi, X, Koiwa, M. (1991 b), Mater. Trans. Jpn. Inst. Met. 32, 48-51. Yasuda, H., Takasugi, T, Koiwa, M. (1992), Acta Metall. Mater. 40, 381-387. Yasuda, K. (1991), in: Concise Encyclopedia of Medical and Dental Materials: Williams, D., Cahn, R. W, Bever, M. B. (Eds.). Oxford: Pergamon, pp. 197-205. Yavari, A. R., Baro, M. D., Fillion, G., Surinach, S., Gialanella, S., Clavaguera-Mora, M. T, Desre, P., Cahn, R. W. (1991), in: High Temperature Ordered Intermetallic Alloys IV: Johnson, L. A., Pope, D. P., Stiegler, J. O. (Eds.). Pittsburgh, PA: MRS, pp. 81-86. Ye, H. Q., Li, D. X., Kuo, K. H. (1985), Phil. Mag. A 51, 829-837. Yegorushkin, V. Y, Kulmentyev, A. I., Rubin, P. E., (1985), Phys. Met. Metall. 60(3), 1-7. Yonezu, I., Fujitani, S., Furukawa, A., Nasako, K., Yonesaki, X, Saito, T, Furukawa, N. (1991), /. Less-Common Met. 168, 201-209. Yoo, M. H. (1991), in: Proc. Int. Symp. on Intermetallic Compounds - Structure and Mechanical Properties (JIMIS-6): Izumi, O. (Ed.). Sendai: The Japan Institute of Metals, pp. 11-20. Yoo, M. H., Fu, C. L. (1991), ISIJInt. 31,1049-1062. Yoo, M. H., Fu, C. L. (1992), Mater. Sci. Eng. A153, 470-478. Yoo, M. H., Fu, C. L. (1993), in: Structural Intermetallics: Darolia, R., Lewandowski, J. X, Liu, C. X, Martin, P. L., Miracle, D. B., Nathal, M. V (Eds.). Warrendale, PA: XMS, pp. 283-292. Yoo, M. H., Horton, J. A., Liu, C. X, (1988), Acta Metall. 36, 2935-2946. Yoo, M. H., Xakasugi, X, Hanada, S., Izumi, O. (1990), Mater. Trans. Jpn. Inst. Met. 31, 435-442. Yoo, M. H., Fu, C. L., Lee, J. K. (1991), in: High Temperature Ordered Intermetallic Alloys IV: Johnson, L. A., Pope, D. P., Stiegler, J. O. (Eds.). Pittsburgh, PA: MRS, pp. 545-554. Yoshida, M., Xakasugi, X, (1991a), in: Proc. Int. Symp. on Intermetallic Compounds - Structure and Mechanical Properties (JIMIS-6): Izumi, O. (Ed.). Sendai: Xhe Japan Institute of Metals, pp. 403-407. Yoshida, M., Takasugi, X (1991 b), in: High Temperature Ordered Intermetallic Alloys IV: Johnson, L. A., Pope, D. P., Stiegler, J. O. (Eds.). Pittsburgh, PA: MRS, pp. 273-278. Yoshida, M., Takasugi, T. (1992), Phil. Mag. A65, 41-52. Yoshihara, M., Suzuki, X, Tanaka, R. (1991), ISIJ Int. 31, 1201-1206.

Yoshimi, K., Hanada, S. (1993), in: Structural Intermetallics: Darolia, R., Lewandowski, J. I, Liu, C. X, Martin, P. L., Miracle, D. B., Nathal, M. V (Eds.). Warrendale, PA: XMS, pp. 475-482. Zakrevskiy, I. G., Kokorin, V V, Shevchenko, A. D. (1986), Phys. Met. Metall. 61 (2), 199-202. Zeng, K., Hamalainen, M., Lilius, K. (1993), CALPHAD17, 101-107. Zeumer, B., Sauthoff, G. (1992), paper at the DGM Annual Meeting in Hamburg. Zeumer, B., Von Keitz, A., Sauthoff, G. (1991), in: 2. Symposium Materialforschung des BMFT: Jiilich, KFA-PLR, pp. 2397-2399. Zhang, M.-X., Hsieh, K.-C, Chang, Y A. (1992a), Mater. Res. Soc. Symp. Proc. 273, 103-112. Zhang, M.-X., Hsieh, K.-C, DeKock, X, Chang, Y A. (1992 b), Scr. Metall. Mater. 27, 1361-1366. Zhang, X, Li, Y, Zheng, Z., Zhu, Y (1991), in: High Temperature Ordered Intermetallic Alloys IV: Johnson, L. A., Pope, D. P., Stiegler, J. O. (Eds.). Pittsburgh, PA: MRS, pp. 137-142. Zhang, W, Chen, G., Wang, Y, Sun, Z. (1993), Scr. Metall. Mater. 28, 1113-1118. Zhang, Y, Lin, D. (1993), Mater. Res. Soc. Symp. Proc. 288, 611-616. Zhang, Y, Xonn, S. C , Crimp, M. A. (1993), Mater. Res. Soc. Symp. Proc. 288, 379-384. Zhang, Y. G., Xu, Q., Chen, C. Q. (1991), in: Proc. Int. Symp. on Intermetallic Compounds - Structure and Mechanical Properties (JIMIS-6): Izumi, O. (Ed.). Sendai: Xhe Japan Institute of Metals, pp. 39-43. Zhao, L., Baker, I., George, E. P. (1993), Mater. Res. Soc. Symp. Proc. 288, 501-506. Zhou, B., Chou, Y. X, Liu, C. X (1993), Intermetallics 1, 217-255. Zhou, X. W, Hsu, X Y. (1991), Acta Metall. Mater. 59,1041-1044. Zhu, J. H., Wan, X. J., Wu, Y. (1993), Scr. Metall. Mater. 29, 429-432. Zinov'eva, G. P., Romasheva, L. F , Saperov, V. A., Gel'd, P. V (1984), Sov. Phys. - Solid State 26, 2115-2117. Zwicker, U., Pack, D., Blaufelder, C. (1979), Z. Metallkd. 70, 514-521.

General Reading Baker, X, Darolia, R., Whittenberger, J. D., Yoo, M. H. (Eds.) (1993), High-Temperature Ordered Intermetallic Alloys V. Pittsburgh, PA: Materials Research Society. Darolia, R., Lewandowski, J. X, Liu, C. X, Martin, P. L., Miracle, D. B., Nathal, M. V (Eds.) (1993), Structural Intermetallics. Warrendale, PA: TMS. Dimiduk, D. M., Miracle, D. B., Ward, C. H. (1992), Mater. Sci. Technol 8, 367-375.

11.14 References

Engell, H.-X, von Keitz, A., Sauthoff, G. (1991), in: Advanced Structural and Functional Materials: Bunk, W. (Ed.). Berlin: Springer-Verlag, pp. 9 1 132. Fleischer, R. L. (1991), in: Proc. Int. Symp. Intermetallic Compounds - Structure and Mechanical Properties (JIMIS-6): Izumi, O. (Ed.). Sendai: The Japan Institute of Metals, pp. 157-164. Izumi, O. (1989), Mater. Trans., JIM 30, 627-638. Liu, C. T., Stiegler, J. O., Froes, F H. (1990), in: Metals Handbook, Vol. 2: Properties and Selection: Non-Ferrous Alloys and Special Purpose Materials. Materials Park: ASM, pp. 913-942. Liu, C. T., Cahn, R. W, Sauthoff, G. (Eds.) (1992), Ordered Intermetallics - Physical Metallurgy and Mechanical Behaviour. Dordrecht: Kluwer. Materials Science and Technology: A Comprehensive Treatment: Cahn, R. W, Haasen, P., Kramer, E. X

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