Resistivity and strength of precipitation-hardened dilute AlAu alloys

Materials Science and Engineering, 48 (1981) 41 - 51

41

Resistivity and Strength of Precipitation-hardened Dilute AI-Au Alloys K. T. HARTWIG, JR., and F. J. WORZALA

Engineering Experiment Station, University of Wisconsin-Madison, Madison, WI 53706 (U.S.A.) (Received May 12, 1980; in revised form September 24, 1980)

SUMMARY

High purity aluminum is often favored over copper for superconductor stabilizer applications because of a lower resistivity and a lower magnetoresistance effect in high magnetic fields. This preference coupled with the major disadvantage of aluminum, namely its lower strength, led to the study reported here on dilute A I - A u alloys. This alloy is strengthened dramatically by a precipitation heat treatment and exhibits low resistivity in the aged condition. Experiments are reported on the yield strength and residual resistivity of dilute A I - A u alloys. Several schemes are developed that allow the conductor design engineer to trade off resistivity and strength. A method of optimizing alloy strengthresistivity is presented.

1. INTRODUCTION

The fact that copper is an excellent cofabrication partner with both NbTi and NbsSn has pre-empted the use of high purity aluminum for similar applications. High purity aluminum is simply too weak to coreduce effectively with NbTi during extrusion and wire drawing and does not possess the diffusion characteristics that would allow aluminum to act like copper during the formation of A15 compounds. The inherent nature of aluminum prevents the formation of A15 superconductor alloys by solid state diffusion and no amount of alloying will remedy this problem. The coreduction drawbacks of high purity aluminum can be altered, however, by alloying. The task is simply to find a solute species that is an effective strengthener, will force rapid and strong work hardening in the alloy and can be rendered ineffective as an electron scatterer in the alloy. Dilute A1-Au 0025-5416/81/0000-0000/$02.50

has the potential for such behavior and was thus chosen for an extensive study of alloy strength and residual resistivity. A study of the basic properties of a carefully chosen alloy could be used to extrapolate predictions for other systems. Some work has been done on the development of highly conductive aluminum alloys. Simoneau and Begin [ 1 ] studied directionally solidified high purity A1-A13Ni in the composition range 2 - 28 wt.% Ni and found that a very pure aluminum matrix could be achieved for alloys of low nickel content. A residual resistivity of 10 n~2 cm was obtained for alloys containing 2.5 wt.% Ni. In a different study Rohatgi and Prabhakar [2] prepared A1-Ni wire 1.78 mm in diameter from cast ingots 41 mm in diameter using either hot extrusion or hot rolling followed by cold drawing. The processing sequence developed led to a general alignment of fibrous dispersed A13Ni without fiber fragmentation or coarsening or fiber matrix separation. In wire form this material possesses an excellent combination of high strength and high room temperature electrical conductivity. Yield strengths range from 155 MPa (at 0.6 wt.% Ni) to 214 MPa (at 6.1 wt.% Ni) in combination with relatively high conductivity. Preliminary work with precipitationhardened A1-Au was begun at the Oak Ridge National Laboratory (ORNL) in 1972 [3, 4]. ORNL was interested in high purity dilute A1-Au alloys for stabilizing superconducting transmission lines. The basis for the study was that gold has an extremely low solid solubility in zone-refined aluminum (perhaps as low as 1 at. ppm at 100 °C) and that remarkable age hardening can be accomplished in dilute A1Au alloys [5]. Solute species cost was not of concern since at the time of the study the price of gold was fixed at U.S. $35 oz-1 and © Elsevier Sequoia/Printed in The Netherlands

42 was not as highly priced on a free floating market as it is today. In the ORNL study [4], high purity alloys containing up to 0.26 wt.% Au were prepared by cold rolling and were subsequently heat treated by homogenization and artificial aging. The residual resistivity was found to decrease dramatically in aged material. The highest value of the residual resist4vity ratio (RRR) R273 K/R4.2 K measured was 2675 and this was for A1-0.12wt.%Au (A1-0.12Au) aged 1000 h at 300 °C. An important conclusion of the study was that aluminum alloys of high RRR are manufacturable. It is to be emphasized that the A1-Au system was chosen for the present study for technical reasons. The advantage of dilute A1Au is that large changes in strength and resistivity are possible through simple heat treatment and such variations allow great flexibility in optimization. Economics was not a part of the present study.

2. BACKGROUND ON PRECIPITATION IN

AI-Au

ALLOYS The phenomenon of age hardening or precipitation hardening occurs in many aluminum alloy systems. Two prime examples are A1-2024 and A1-6061. The 2000 series alloys contain copper as the primary alloying element while the 6000 series alloys contain magnesium and silicon as primary solutes. In both alloy series the solute species can be heat treated to form tiny insoluble crystallites within the aluminum matrix. The crystallites or precipitates hinder matrix deformation a n d thus act as strengtheners. In order to accomplish effective age hardening it is necessary that there be considerable solute solid solubility at elevated temperatures and that solubility decrease with decreasing temperature. If these conditions exist, then precipitation strengthening can be induced by heat treatment. Such heat treatment includes a solutionizing step, a quench and an aging step. Solutionizing is a baking operation carried out at a temperature within the solid solution phase region and is intended to homogenize alloy composition on a microscopic scale and to place the maximum amount of solute into solution. The quench, a rapid cooling to low temperature, follows

solutionizing immediately. Aging follows quenching and is the moderate temperature (often 200 - 300 °C) heat treatment necessary to drive the solute out of solution and into pure solute, solute-rich or intermetallic compound clusters or precipitates. Initially, the strength of the alloys is that of the supersaturated solid solution. If the first precipitates to form after aging are small and coherent with the matrix, they will be cut by moving dislocations and the alloy will not be much stronger than the initial material. As precipitates grow with increased aging and possibly change their internal structure or the nature of their interface with the matrix, the work required to cut each particle may increase and the alloy strength will increase. At some point, the alloy strength will be a maximum and is dependent on the yield stress of the solid solution, the volume fraction of precipitates, the precipitate size and spacing and the precipitate-matrix interface energy. As aging progresses still further, precipitates will become larger and more widely spaced, and moving dislocations formed as a consequence of deformation will be forced around particles instead of through them. Further increases in particle size and separation distance lead to a decrease in yield strength (the overaged condition). In 1967, von Heimendahl [5] studied the solubility of gold in aluminum and found that the solid solubility reaches a maximum of about 0.3 wt.% Au at the eutectic temperature of 640 °C and decreases with decreasing temperature. In a later study, Fujikawa and Hirano [6] accurately determined the location of the solid solution phase boundary and their result is reproduced in Fig. 1. yon Heimendahl reported remarkable age hardening in the alloy as shown in Fig. 2. Sankaran [7] and Sankaran and Laird [8] studied the growth kinetics and precipitate morphology in dilute A1-Au by electron microscopy. They found that aging a supersaturated solid solution of aluminum containing 0.2 wt.% Au (A1-0.2Au) produces a fine distribution of plate-like precipitates of the equilibrium phase ~-A12Au on matrix cube planes, without the formation of any metastable phases as suggested by Fujikawa e t al. [12]. Sankaran and Laird suggested that the plates nucleate on dislocation loops formed by the condensation of vacancies which are

43 GOLD (ATOMIC %) 0 0 2 0 0 3 0.04 0 0 5

001 800

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CHAJ~CTERI/ATION

006 l

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FABRICATION SEQUENCE

L

A1

700 I

600

INDUCTIONMELT IN GRAPHITE UNDER ARGON

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D 500

1 / w

P- 400

COMPOSITIONANALYSISOFAL&AU. ] ELECTRICALRESISTIVITYMEASUREMENTSONAL AT 273, 77 AND4.2 K

Au

MACHINE

COMPOSITION ANALYSIS METALLOGRAPHY

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05

HOMOGENIZATIONAT 600 "C FOR12 HRANDQUENCH INTO15"CWATER

Fig. i. The aluminum-rich side of the AI-Au phase diagram [ 7 ].

I ELECTRICALRESISTIVITYMEASURE-I MENTSAT 273, 77 AND4.2 K

I I SOTHERVALHEAT TREATMENTIN OIL ORSALTBATH 0 < T < 350"C 0 < t < 300 HR

9ooo

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I METALLOGHAPHY HARDNESSMEASUREMENTS TENSILE MEASUREMENTSAT 296, /7 AND 4.2 K

I ELECTRICALRESISTIVITYMEASUREMENTSAT 273, 77 AND4.2 K ELECTRONMICROSCOPY

rooo

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Fig. 2. Variation in yield strength with aging time for A1-0.2Au [5].

retained during quenching. They found a matrix-precipitate lattice mismatch of 4.8% parallel to the broad faces of the precipitate platelets and further suggested that this mismatch produces a large coherency strain field which gives rise to the large age-hardening effect. Misfit dislocations develop at longer aging times, accommodate the mismatch and lead to softening. 3. EXPERIMENTAL PROCEDURES

The experimental procedures used for sample preparation, analysis and testing are outlined in Fig. 3 and presented in detail elsewhere [ 9]. Alloy preparation involved induction melting 99.999% A1 with 99.999% Au in high purity graphite crucibles under argon. Ingots 25.4 mm in diameter containing 0.0, 0.5, 0.1, 0.15 and 0.2 wt.% Au were extruded at 525 °C to diameters of 6.4 and 1.6 mm.

All samples were homogenized in air at 600 °C for at least 2 h and quenched rapidly into cold water. Aging heat treatments were performed in an isothermal bath of oil or liquid salt at temperatures between 200 and 350 °C. Neutron activation analysis was used to determine accurately the gold content of each specimen. The alloy impurity content was determined by emission spectroscopy to be 0.001 wt.% Ag, 0.001 wt.% Cu, 0.0001 wt.% Fe, 0.0001 wt.% Mg and 0.0001 wt% Si. Specimens at least 50 mm long and 6.4 mm in diameter were used for resistivity measurements. Conventional potential difference data were taken at 273, 77 and 4.2 K while an eddy current decay method was also used for measurements at 4.2 K. Uniaxial tensile tests were performed on chemically cleaned specimens 5.3 mm in diameter with an Instron machine at a strain rate of 8.4 × 10 -4 s-1. A clip-on 24.5 mm extensometer was attached to the specimens to monitor strain. The determination of the yield strength o0.2 at 0.002 permanent strain offset was made without correction for the reduction in cross-sectional area accompanying strain. Measurements were taken at room temperature in air, at 77 K under liquid nitrogen and at 4.2 K under liquid helium.

44

The temperature dependence of yield strength was determined by single-specimen measurements of o0.9 at different temperatures following a method similar to that of Cottrel and Stokes [10].

4. EXPERIMENTAL RESULTS AND DISCUSSION

4.1. Electron microscopy Figure 4 is a transmission electron photomicrograph of A1-0.2Au aged for 1 h at 250 °C after solution heat treatment. Such an aging treatment leads to an alloy of moderate strength. Plate-like A12Au precipitates with an average length of 250 £ and a thickness of less than 30 k are readily visible in the photomicrograph. The average separation and density are 800 A and 1.9 × 1015 cm -3 respectively. Precipitates are spaced uniformly within the crystal grains and are oriented on {100} habit planes. In deformed areas of foil specimens it was noted that the small particles of A12Au indeed act as barriers and are not easily cut by moving dislocations.

Fig. 4. Transmission electron photomicrograph of AI-0.2Au aged for 1 h at 250 °C (peak aged). A12Au precipitates are visible.

Figure 5 is a transmission electron photomicrograph of A1-0.2Au aged for 30 min at 300 °C. In this overaged condition the specimen exhibits a microstructure of oriented and finely dispersed precipitation platelets varying from 250 to 1400 ~ in width within a high purity matrix. The particle density and separation are roughly 8 × 1014 cm - s and 1000 A respectively.

Fig. 5. Transmission photomicrograph of AI-0.2Au aged for 30 min at 300 °C (overaged). A12Au precipitates are visible.

4.2. Resistivity The addition of gold to A1-Au alloys increases resistivity in direct proportion to the quantity of gold in solid solution. Gold content also affects the precipitation kinetics and the final level of resistivity attained as a consequence of precipitation (Table 1). For concentrations less than about 0.015 at.% Au (0.13 wt.% Au), a linear relationship defined by Ap = 2.1 + 0.1 p~2 cm (at.%) -1 is measured (Fig. 6) and is in good agreement with the value of 2.0 ~ 2 cm (at.%) -1 reported by Fujikawa and Hirano [6]. For concentrations greater than 0.015 at.% Au, complete solubility of the gold was not achieved by homogenization at 600 °C. This is shown by the departure from linearity of the curves in Fig. 6. The intrinsic resistivity, or phonon contribution, is the major contributor to resistivity at 273 K. At 4.2 K, however, the temperature contribution drops to less than 1% of the total and the contribution from the solute dominates. Because contributions to the resistivity from such factors as size effect, dislocations, stacking faults, grain boundaries and base metal impurities are estimated to be less than 10% of the total [9], it is reasonable to assume that the factor contributing the most to residual resistivity P0 is the dissolved solute. Since the gold-bearing precipitates (A12Au) in the aged alloy are relatively far apart and their strain fields are localized, they do not contribute significantly to the residual resistivity. The precipitate contribution to resistivity probably results primarily from

45 TABLE 1 Resistivity of aluminum and AI-Au alloys Temperature and alloy condition

(n~ cm)

Resistivity 0 wt.%Au

0.064 +- 0.004 wt.% A u

0.10 +- 0.005 wt.% A u

0.15 +- 0.01 wt.% A u

0.21 +- 0.01 wt.% A u

2450 2440

2460 2450

2470 2460

2480 2470

2490 2480

4.0 0.5

16.5

28 3.2

37 4.0

41 5.0

273 K

Homogenized Aged for 10 h at 300 °C 4.2K

Homogenized Aged 70

I

I

14 s //

6O ¢1 c

tiny change in resistivity measured at room temperature. The resistivity of solution-heattreated A1-0.2Au is 2490 n~2 cm at room temperature and only 41 n~2 cm at 4.2 K. The corresponding values for overaged material are 2480 n~2 cm and 5.0 n~2 cm. The principal results of aging dilute A I - A u are as follows. (1) Aging progresses slowly at ambient temperatures b u t always results in a decrease in resistivity. (2) The absolute value of resistivity is directly proportional to the quantity of gold in the alloy irrespective of heat treatment. (3) The decrease in resistivity with aging time is independent of the quantity of gold in the alloy. (4) The shape of the aging curve is independent of composition for a given aging temperature. (5) An increase in annealing temperature accelerates the aging process.

t

~,at % Au/

50

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Fig. 6. Dependence of resistivity at 4.2 K on gold concentration for solution-treated AI-Au.

volume and size effects and both of these are small. The effect of aging on the residual resistivity is shown in Fig. 7. The dramatic variation in P0 with aging contrasts sharply with the Io z # / '

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46

(6) After the initial major drop in resistivity takes place, a much slower steady decrease with time is observed. The last effect (point (6)) involves the slow growth of precipitates by absorption of gold atoms from the surrounding solid solution after the diffusion fields of adjacent precipitates have overlapped. If the decrease in resistivity with increased aging time for A1-Au is independent of gold content (point (3)), then for a given aging temperature pi/ci should be constant where Pi is the resistivity for an alloy of composition ci. From Fig. 8 it can be seen that resistivity is proportional to gold content irrespective of heat treatment. The logarithm of residual resistivity divided by the atoinic fraction of gold is plotted against the logarithm of aging time for aging at 200, 250 and 300 °C. A decrease with time independent of gold concentration exists for all temperatures. The curves are drawn through data for A1-0.15Au; the vertical error bars bound scatter in the ratio p~/ci for the composition range 0.1 - 0.2 wt.% Au. If it is assumed that the largest contribution to residual resistivity is from the dissolved gold, then the curves in Fig. 8 indicate that the fraction of original solute removed by aging is independent of alloy composition. This result is in qualitative agreement with work by Turnbull [11] on A1-Ag and A1-Cu alloys.

4.3. Yield strength Uniaxial tests at temperatures of 295, 77 and 4.2 K were performed on specimens 5.3 mm in diameter of pure aluminum and aluminum containing nominally 0.10 and

0.20 wt.% Au. The pure aluminum samples were given a strain relief anneal at 525 °C for I h in air before testing. Samples of the dilute alloy were tested in the solution-treated peakaged (1 h at 250 °C) condition and in the overaged (0.5 h at 300 °C) condition. The specimens of pure aluminum had an RRR of between 3500 and 4000 and an equiaxial grain structure with an average grain diameter of 0.07 mm. The average yield strength of the pure aluminum was found to be 5.5 + 0.6 MPa at room temperature, 6.9 + 0.7 MPa at 77 K and 9.0 + 1.4 MPa at 4.2 K. Uncertainty in these values represents one standard deviation for the several measurements taken. Analysis of the data from specimens strained at more than one temperature led to the strength ratios presented in Table 2. These ratios are the average of several tests and are reproducible to within 5%. The strength of pure aluminum at 4.2 K is noted to be 48% above that at room temperature. Average yield strength values for alloys are also listed in Table 2. The aging effect is seen to be greatest at 250 °C and in A1-0.2Au where a peak strength of 43.4 MPa is measured. A summary of the effects of the atomic fraction of gold, the temperature and the aging conditions of temperature and time on the yield strength at 0.002 strain for dilute A1-Au is given in Table 3. The strength dependence comments are derived from tests on multicrystal specimens, where in many cases the number of grains per gauge cross section was less than eight and the average grain diameter d was 2.5 mm. Grain size effects

E a" .1 0 ~i I0

F-

AGING TIME (hr)

Fig. 8. D e p e n d e n c e o f resistivity at 4.2 K per atomic fraction o f gold on aging t i m e for A I - A u : I, composition range, 0.1 - 0,2 wt.% Au.

47 TABLE 2 Temperature dependence of yield strength for pure aluminum and dilute A1-Au Yield strength at 2 9 5 K (MPa)

Pure a l u m i n u m Annealed

S t r e n g t h ratio

00. 2 at 4 . 2 K

00. 2 at 77 K

00. 2 at 4.2 K

Oo.2 at 2 9 5 K

00. 2 at 2 9 5 K

00. 2 at 77 K

5.5

1.48

1.28

1.15

Al-O. 1 A u Solution heat treated Peak aged Overaged

12.8 26.2 17.9

1.29 1.25

1.19 1.19

1.09 1.06

Al-O.2Au Solution heat treated Peak aged Overaged

13.4 43.4 26.5

1.83 1.27 1.25

1.37 1.18 1.18

1.33 1.07 1.06

TABLE 3 Strengthening effects in dilute A1-Au Variable

Atomic fraction c of gold Solid solution alloy Aged alloy Measurement temperature T M Solid solution alloy

Aged alloy Aging conditions Time

Temperature

Strength dependence at 0 . 0 0 2 strain

c 112 c

Alloy is 1.5 - 2 times stronger at 4.2 K than at 300 K TM2/3 for T M < 295 K Strength increases to a maximum and then drops For higher temperatures the maximum strength and the time to maximum strength decrease

were not considered, although specimens with a finer grain size showed higher yield strengths. The grain size in the alloy will contribute to strength in the same fashion as it does in pure metal, i.e. by way of an additive d -112 dependence.

4.4. Stabilizer property optimization It is necessary to assess the physical, manufacturing and economic characteristics of a

particular material considered for use as a superconductor stabilizer. All the factors mentioned are important, but for a specific application one or two particular characteristics m a y be overriding. For a large superconductive magnetic energy storage system, cost and mass are most important. Low density and low resistivity are preferred. Strength is also an attractive characteristic since large magnetic forces are generated within any magnet conductor. Because an increase in stabilizer strength for a particular alloy type is accompanied by a resistivity increase, the use of a stronger stabilizer material will reduce the a m o u n t of c o n d u c t o r structural material required for a particular application but will increase the a m o u n t of stabilizer material required for conduction purposes. The overall effect on c o n d u c t o r cost is therefore difficult to assess. For aluminum or a dilute aluminum alloy as a stabilizer, the material density will n o t affect cost because density is essentially the same for both materials. The quantity of alloying elements added and any special thermal-mechanical treatment required, however, will affect cost, especially when the addition is very expensive and the alloy processing schedule is complicated. These factors are difficult to quantify; a scheme for the assessment of the alloy stabilizer potential that takes into account the alloying element costs and processing costs has n o t been formulated.

48

Since the residual resistivity P o and the yield strength 00.2 are easily quantified and are both important to the stabilizer, a tradeoff and optimization scheme involving both Oo.2 and P0 at 4.2 K is worthwhile. General consideration should first be given to a comparison of 00.2 and Po at various aging temperatures and times. For a particular alloy composition a "resistivity-strength map" can be prepared that provides a scheme for comparison involving all four parameters. Figure 9 presents isoresistivity and isostrength curves in a field of aging temperature and aging time. The data are from measurements taken at 4.2 K on age-hardened A1-0.2Au specimens 6.4 mm in diameter. Similar maps are easily derived for more dilute alloys since the composition dependences are known. To a first approximation, the quantity of gold present in the hardened alloy is directly proportional to the increase in yield strength and inversely proportional to the residual resistivity. 400

I

I

I

I

values of residual resistivity can be attained in A1-0.2Au aged to a particular yield strength and at which temperature the alloy should be heat treated to acquire the specific yield strength or residual resistivity desired. For optimum design, Fig. 10 is best. Yield strength at 4.2 K is plotted against residual resistivity for aging at 150, 200, 250 and 300 °C. The 150 and 200 °C aging curves are plots derived from yon Heimendahl's yield strength results. The locus of all points defining the maximum strength for a given resistivity is approximated by the broken curve in Fig. 10 labelled "estimate of boundary limit". This estimate curve is for material of maximum yield strength and minimum resistivity that can be formed by simple heat treatment of A1-0.2Au. The aging temperature and time necessary to form an alloy of optimum conditions can be selected by using either yield strength or residual resistivity aging curves in conjunction with the estimate curve. The value of relating yield strength directly to residual resistivity without considering aging conditions becomes immediately apparent. 8O

JSORESISTIVITY CONTOURS (n~Q~m)

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150

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I0 TIME

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Fig. 9. Resistivity-strength map for A1-0.2Au at 4.2 K. Isoresistivity contours and isostrength contours are plotted in a field of aging temperature and aging time. The peak yield strength at 0.002 strain curve is shown.

Figure 9 overlaps P0--TA--tA and a0.2-TAt A plots where T A is the aging temperature is the aging time, so that a direct comparison may be made between P o and 00.2 for various heat treatment conditions. A magnet designer can determine quickly which and t A

Fig. 10. Dependence of yield strength at 4.2 K on 4.2 K for AI-0.2Au. The curve labelled "estimate of boundary limit" (- - - ) defines the boundary beyond which strength cannot exceed for a

resistivity at

given resistivity. Another way to assess the importance of changes in properties with aging is to define a figure of merit that presents the pertinent information in a compact form so that easy comparison of alternative heat treatment techniques is possible. With dilute AI-Au, it is appropriate to consider ~o/Ap where Aa is defined as the increase oa,oy -- apu~e~ in ao.2 at 4.2 K caused by alloying and heat treating and where ~P0 is the corresponding increase P a.oy -- P ~re Ai in residual resistivity for the

49

alloy. Large yield strengths and small resistivities are attractive; high values of the ratio of A o / A p o are desirable. Figure 11 shows curves of A o / A p plotted against aging time at heat treatment temperatures of 200, 250 and 300 °C for the A10.2Au alloy. For all cases the ratio increases with time until well after peak hardness is achieved. The curves exhibit the following trends. (1) Early stages of aging involve a rapid increase in A o/Ap o. (2) After peak yield strength the rate of increase in A o / A p o slows down dramatically. The A o / A p o versus tA curve may show a first maximum or a plateau shortly after peak yield strength is reached. (3) As aging continues, A o / A p o may again rise and continue to do so until the matrix is devoid of all stray gold atoms. A constant A o is reached before a constant Ap 0 is reached.

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-o

y

o

--r~

~ PEAKYIELDSTRENGTH

I

~'

I

o

I

,o'

I

~o'

I

,o~

TIME (hr)

Fig. 11. Increase m yield strength at 4.2 K per increase in resistivity, Ao/Ap vs. aging time for A10.2Au. The location of peak yield strength is shown.

Figure 11 is useful for comparing the figure of merit, A o / A p o , at a given time for alloys aged at various temperatures. It shows that one heat treatment temperature may be better than another for a given aging time from the standpoint of increased strength and decreased resistivity. Since strength and resistivity are often not of equal importance, the value of Fig. 11 for engineering design may be less than those of Figs. 9 and 10.

4.5. Comparison w i t h o t h e r materials

The ultimate test of potential for the use of dilute A1-Au or any other metal or alloy for a stabilizer is direct comparison of pertinent property data with alternative choice materials. Since many properties must be considered, comparison is not straightforward. The problem is compounded because it is difficult to assign quantitative values to specific properties in many cases. To simplify comparison, only yield strength and residual resistivity are discussed. In Table 4 information is listed for comparison. A variety of aluminum and copper pure metals and alloys that provide a wide range of strength and resistivity values are included. Also listed for each material is a figure of merit, ao.2/p0, similar to the A o / A p o term discussed above. The higher the figure of merit, the more attractive is the material, a0.2/po (×10 ~5 Pa ~ - 1 cm-1) varies from a low value of 0.3 for Al-ll00(0) to a high value of 409 for oxygen-annealed high purity copper. It is clear that alloys and impure metals such as AI-ll00 have relatively low figures of merit while pure metals have high figures of merit. Specialty alloys such as A1-Ni and A1-Au can have high figures of merit if given the proper thermal-mechanical treatment. Such a treatment results in the formation of a stable second-phase intermetallic compound within a relatively pure metal matrix. Figures of merit for A1-Au and A1-Ni can be as high as 6.6 and 2.5 respectively and compare favorably with merit figures for aluminum of RRR 3500 and for OFHC copper of RRR 100. Yield strength alone or residual resistivity alone is often more important for stabilizer design than an equal combination of both. It is therefore worthwhile to present yield strength and resistivity as shown in Fig. 12 where the yield strength at 4.2 K is plotted against the residual resistivity for a variety of materials. Figure of merit contours for Oo.~/p o values of 1017, 10 is and 1014 are drawn in the figure. One general feature is apparent: as resistivity decreases, so does the strength for a particular metal or alloy. The trend, however, is not always followed, as evidenced by directionally solidified A1-6.1Ni and oxygenannealed copper. Both materials have extremely high yield strengths for their associated residual resistivities. The drawback of

50 TABLE 4 Yield strength at 4.2 K and residual resistivity for aluminum and copper conductors

Alloy

Condition

Reference

00. 2 (MPa)

Po (n~ era)

Oo.2/Po (xlO 15 Pa ~ - 1 cm-1)

99.999A1 Al-ll00 AI-1350 A1-5005 A1-6063 Al-6201 AI-0.2Au

Annealed Annealed Annealed Cold drawn Aged T5 Aged Peak aged Overaged Annealed Annealed Annealed Cold drawn Annealed Oxygen annealed

PS [13] o [ 14] ;p, PS o [14] ;p, PS a [ 1 5 ] ; p [16] a [14] ;p, PS PS PS o [14] ;p, PS o [14] ;p, PS o [14] ;p, PS PS [13] o, estimate; p [17]

9.0 a 65 at 20 K 100 at RT 320 at RT 165 at 77 K 347 at RT 57 33 83 at RT 87 at RT 99 at RT 187 at RT 90 at 20 K 90

0.7 200 73 460 280 350 17 5 70 60 71 76 16 0.2

12.9 0.3 (1.4) b (0.7) 0.6 (1.0) 3.3 6.6 (1.2) (1.5) (1.4) (2.5) 5.6 409

AI-0.3Co-0.5Ni (NiCo) AI-0.5Co-0.5Fe (Super-T) Al-0.55Fe (Triple E) A1-6.1Ni OFHC Cu 99.999Cu

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with a r o o m t e m p e r a t u r e yield strength above 6 9 MPa a n d a figure o f m e r i t in excess o f 7 × 1015 Pa gZ-1 c m -1. I f an a l u m i n u m a l l o y c o u l d easily b e f a b r i c a t e d w i t h t h e s e p r o p erties, it w o u l d b e m u c h m o r e a t t r a c t i v e t h a n O F H C c o p p e r f r o m the standpoints of yield s t r e n g t h , resistivity, w o r k a b i l i t y a n d p r o b a b l y c o s t f o r large stabilizer a p p l i c a t i o n s . T h e t r e n d e x h i b i t e d in Fig. 12 i n d i c a t e s t h a t t h e f o r m a t i o n o f such an a l l o y is possible.

51 5. C O N C L U S I O N S

High purity aluminum can be combined with a small a m o u n t of gold to form a relatively strong low resistivity alloy. These properties are the result of a precipitation heat treatment that causes gold in solid solution to precipitate in the form of A12Au platelets that are effective in blocking dislocation motion and contribute less to bulk resistivity than individual gold atoms would if dispersed homogeneously throughout the aluminum matrix. To a first approximation the increase in yield strength is directly proportional to the quantity of gold present for the aged alloy and proportional to the square root of the concentration for the solid solution alloy. The residual resistivity always decreases with aging and is directly proportional to the quantity of gold present in the alloy irrespective of heat treatment conditions. Specific results of the present study are as follows. (1) The yield strength of high purity aluminum at 4.2 K is about 50% above that at room temperature. (2) The artificial aging of quenched dilute A1-Au alloys results in a substantial decrease in residual resistivity. Trends in the resistivity -aging time behavior indicate that (a) the mechanism responsible for precipitation is similar to that responsible for the annealing of quenched-in vacancies in pure aluminum, (b) the precipitation mechanism is independent of composition up to at least 0.20 wt.% Au, (c) the residual resistivity is directly proportional to the a m o u n t of gold in the alloy irrespective of heat treatment and (d) even after long aging times the residual resistivity continues to decrease at a slow rate with increased aging. (3) Precipitation-hardened A1-0.2wt.%Au can be eight times as strong as pure aluminum at 4.2 K. The strength of the alloy was found (a) to be directly proportional to the gold concentration c for aged alloy, (b) to be proportional to c z/2 for solution-treated alloys and (c) to have a temperature dependence proportional to T 2/3. (4) The A1-Au alloy yield strength and residual resistivity as functions of aging conditions can be compared conveniently on a resistivity-strength map for alloy property

determinations appropriate for stabilizer design (see Fig. 9). (5) The optimum yield strength-residual resistivity relationship for A1-0.2wt.%Au alloy is established (see Fig. 10). (6) A figure of merit, defined as a0.2/P0, indicates that dilute A1-Au alloy is as good as aluminum with an RRR of 1000 and OFHC copper with an RRR of 100 for stabilizer material. ACKNOWLEDGMENTS

The financial support for this investigation was provided by the National Science Foundation, the U.S. Energy Research and Development Administration, The Wisconsin Electric Utilities Research Foundation and the University of Wisconsin. REFERENCES 1 R. Simoneau and G. Begin, J. Appl. Phys., 44 (1973) 1461. 2 P. K. Rohatgi and K. V. Prabhakar, Metall. Trans. A, 6 (1975) 1003. 3 Cryogenic power transmission technology :-- cryogenic dielectrics, ORNL Q. Rep. TM-5030, September 1975, p. 17 (Oak Ridge National Laboratory). 4 Cryogenic power transmission technology -cryogenic dielectrics, ORNL Semi-annu. Rep. (March 1, 1973) TM-4187, May 1973, p. 53 (Oak Ridge National Laboratory). 5 M. yon Heimendahl, Z. MetaUkd., 58 (1967) 230. 6 S. Fujikawa and K. Hirano, J. Jpn. Inst. Met., 38 (1976) 929. 7 R. Sankaran, Thesis, University of Pennsylvania, 1973. 8 R. Sankaran and C. Laird, Philos. Mag., 29 (1974) 179; Acta Metall., 22 (1974) 957; 24 (1976) 517. 9 K. T. Hartwig, Ph.D. Thesis, University of Wisconsin-Madison, 1977. 10 A. H. Cottrell and R. J. Stokes, Proc. R. Soc. London, Set. A, 233 (1955) 17. 11 D. Turnbull, H. S. Rosenbaum and H. N. Treaftis, Acta Metall., 8 (1960) 277. 12 S. Fujikawa, K. Hirano and M. Hirabayashi, Z. MetaUkd., 59 (1968) 782. 13 Handbook on Materials for Superconducting Machinery, Mechanical, Thermal, Electrical and Magnetic Properties of Structural Materials, MCKHB-04, Metals and Ceramics Information Center, Battelle Columbus Laboratories, Columbus, OH, November 1975. 14 E. H. Chin, personal communication, September 1978. 15 Aluminum Standards and Data, Aluminum Association, 750 Third Avenue, New York, 3rd edn., January 1972. 16 R. L. Powell, W. J. Hall and H. M. Roder, J. Appl. Phys., 31 (1960) 496. 17 F. R. Fickett, Annu. Rep. INCRA Project 186A, August 1973.