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.—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» y6a + 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 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} 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) 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

o o o o

s a posit ions

o N

a o

o "33 o

'a

6 ON

4 (3


in

p q p o p

8S

d> d> o> d> ^ p 2 r i r n .^- > ... : .. r r > > ^ s : ^ ^

WE43

^"*-

"*"

—

Q_

**«

11»

100

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

• •

200-

2000)

03 N \

100

- " WE43

s

^ ^

B

— AE42 13

v

•

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

o &

a a a a a a a a

a

© o o o o o o o

o

CO

CO

CO

CO

CO

CO

CO

CO

© © © © © © © * ©

§

88 es

Figure 4-29. Microstructure of an Mg-Mg 2 Si alloy (courtesy of Metallgesellschaft).

© co © © co © © © © © ©'©'©©© X d

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 ^

X d

X X cti B!

a a a aa 8 2 S©8 o

© " © © © X X

a a ©

©

a a ©

©

©

a a ©3

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.



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.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 ^.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

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
~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 -> 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 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.

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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

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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

278

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

279

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

280

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

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 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

36-48 850 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 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 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 or 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 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 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 K11156K 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 ( 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 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 -482-200°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