On the effect of ordering upon the strength of Cu3Au

ON THE EFFECT OF ORDERING

UPON THE STRENGTH

OF CuJu”

G. W. ARDLEYt Experiments have been carried out to examine the effect of ordering on the mechanical strength of single crystals of CurAu. It was found that as the disorderdorder reaction proceeds the strength is affected in two ways. Firstly, as the degree of long-range order within the anti-phase domains increases the roomtemperature strength decreases. Secondly, as the anti-phase domains grow larger the room-temperature strength first increases and then decreases. At the ordering temperature there is an appreciable drop in strength as the alloy passes from the ordered to the disordered state. Above the ordering temperature the strength increases with increasing temperature to a maximum value and then decreases again. The temperature at which this maximum occurs depends upon the applied strain-rate.

EFFET

DE L’ORDRE

SUR LA RESISTANCE

DE

CurAu

Ces exp6riences avaient pour but d’ttudier l’effet de l’ordre sur la rssistance m&canique de monocristaux de CuaAu. La &action desordre-ordre a deux effets sur la resistance: lo-Quand l’ordre a grande distance augmente, la rdsistance %la temperature ordinaire diminue. 2”-Quand les domaines en antiphase grossissent, la r&&stance croit, puis cl&roit, A la temperature de Curie, ii y a une chute appr&iable de la resistance quand l’alliage passe de l&at ordonn6 $ F&at d&ordonnC, mais au-dessus de cette temperature, la r6sistance crdt avec la temperature, puis d&molt; la temp&rature B laquelle le maximum se produit depend de la vitesse de deformation. EMBER DEN

EINFLuSS DES ORDNUNGSVORGANGES FESTIGKEIT VON CurAu

AUF

DIE

Es wurde der Einfluss des Ordnungsvorganges auf die mechanische Festigkeit von AuCu~-Einkristallen untersucht und dabei festgestellt, dass der Vorgang Unordnung~Ordnung die Festigkeit auf zweierlei Weise beeinflusst: Erstens nimmt die Festigkeit bei Raumtemperatur mit zunehmendem Fernordnungsgrad in den Antiphasen-Bereichen ab. Zweitends nimmt die Festigkeit bei Raumtemperatur zu und flllt dann wiederum ab, wenn die Antiphasen-Bereiche grosser werden. Bei der kritischen Temperatur, also beim ii’bergang vom geordneten in den ungeordneten Zustand, wurde ein betrachtlicher Abfall in der Festigkeit festgestellt. Oberhalb der kritischen Temperatur nimmt die Festigkeit bis zu einem Maximalwert zu, urn dann widerum abzufallen. Die Temperatur, bei der dieser Maximalwert auftritt, hgngt von der Grosse der angelegten Spannung ab.

PREPARATION

INTRODUCTION

eneral

Electric

22, 1954. Research Laboratory,

Schenectady,

New

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CRYSTALS

The alloys were made by melting together known amounts of high purity (99.999~o) copper and gold to give the CuBAu composition. The copper and gold were first cut into small pieces measuring approximately 1/8X 1/8X .040 in., thoroughly mixed, and then melted together in a graphite crucible under a pressure of less than 10V5 mm Hg. Prior to this the crucible had been cleaned and degassed by heating it to 1300°C for two hours under a high vacuum. Induction heating was used so that the stirring action would help to mix the constituents. The alloy was held molten for about 15 minutes and then allowed to cool to room temperature. To be sure of an homogeneous alloy this procedurei.e., cutting into small pieces, thoroughly mixing the pieces and then melting them together-was repeated twice more on each ingot before it was cold-swaged down to .lOO in. wire. The crystals were grown by the Bridgman method using a lowering rate of 1” per hour. A split graphite mould was used and a pressure of 10-h mm Hg was maintained above the crystals throughout the growth period. Square crystals were preferred for these experiments and two sizes were used, 1.5 mm and 2.5 mm on a side.

It has been known for some time now that the mechanical strength of an ordering alloy changes as the alloy passes from the disordered to the ordered state. It was thought that in alloys of the Cu& type the ordered structure was always weaker than the disordered one, whilst in alloys of the CuAu type the reverse was truel+ CL&U remains face-centered-cubic upon ordering but CuAu changes from f.c.c. to f.c. tetragonal. This led to the belief that a change in crystal structure upon ordering was necessary to produce an increase in strength .5*6,7In this paper we shall consider only the case of CuaAu and it will be shown that one can obtain an increase in strength upon ordering afthough it is not a permanent increase but rather a transitional one similar to that observed in CusAu by Broom and Briggs,R and in CuAu by Dehlinger and Crafg and Nowack.’ A complete review of the literature up to 1940 on “age-hardening” in precious metal alloys, together with an extensive bibliography, can be found in the paper by Vines and Wise.l* ’ Received December

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To hold a crystal in the tensile testing machine, a steel grip was brazed to each end of the crystal. This was done by having a U-shaped grip that allowed the crystal to lie between its prongs with only about .OOl in. clearance on either side. The crystal and grips were assembled on a flat quartz plate, and a piece of brazing solder was placed at each grip. The whole assembly was put into a horizontal quartz tube and heated to 780°C under a protective atmosphere of hydrogen, whereupon the brazing solder melted and ran down into the gaps between the grips and the crystal, forming a good joint. The solder also ran out along the crystal a little way and formed a meniscus-shaped fillet between the specimen and the end of the grip. This is advantageous because it avoids any sharp discontinuity in cross-section in going from specimen to grip. After brazing, the specimen cooled to room temperature in just a few minutes. Whenever the specimens were to be annealed for an extended period of time they were sealed into glass tubes under a vacuum, and heated in a resistance furnace. When they were to be annealed for short periods of time they were immersed into a silicone oil bath that was heated by an electric mantle heater. If the surface

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of the oil was protected from the atmosphere by passing nitrogen across it, the oil could be heated to as high as 380°C for several days before any serious thermal decomposition occurred. EXPERIMENTAL

RESULTS

I

Before describing the main experimental results there are a few observations on the mode of deformation in CulAu that are worthy of note. Figure la shows a reproduction of a load elongation curve for a single crystal of CuaAu deformed at room temperature. There are two interesting features to be noticed in this curve : (1) the deformation occurs jerkily at first and then becomes smooth, and (2) the stress does not increase until the jerky flow has ceased. It appears that each drop in load is associated with the formation of a highly localized slip band so that after several of these jerks have occurred, the specimen looks like the one shown in the inset. Further deformation causes more slip bands to appear but from a visual observation they always seem to occur in new parts of the crystal. Within the jerky flow region, therefore, one will always be measuring the strength of undeformed material. The jerky flow is

160

140

120

100 ti 3 f

80

2 s 60

40

20

0I

_ 2

TOP FACE

SIDE

(a)

FIG. 1. Typical

load-elongation

curves

-

showing

FACE

X5

I % -

ELONGATION jerky flow. Inset-slip bands on two faces of a crystal deformation in the jerky range.

1 W

after a small amount

of

ARDLEY:

IIll r IIT

STRENGTH

I

FIG. 2. Room temperature strength rows domain single crystal of CU3Au ordered at 350°C. The region line ABC represents the contribution to the strength anti-phase domain boundaries.

size for a above the from the

usually much more pronounced than that which is shown in Fig. la. The more usual size is like that shown in Fig. lb. The observations described above have been separated from the main body of data because as yet they are only of a very preliminary nature. They do bear a strong resemblance however to the observations of Ardley and Cottrellu who found similar jerky flow characteristics in B-brass. EXPERIMENTAL

RESULTS

II

The data shown in Fig. 2 are typical of that for the room-temperature strength of a single crystal of CusAu as a function of the anti-phase domain size for an annealing temperature of 350°C. These data were obtained in the following way. First, the specimen was obtained in the disordered state at room temperature, by heating it to 420°C for several hours and then quenching it to room temperature in water. The room-temperature resistivity of the specimen was measured potentiometritally and then the specimen was strained in tension to measure its strength. The value obtained for the roomtemperature strength of the disordered alloy is denoted by the point A in Fig. 2. To vary the domain size, the specimen was heated to the desired temperature, 350°C in this case, in an oil bath. The specimen was put into the bath for a short period of time and then it was quenched to room temperature in order to measure its room-temperature resistivity and room-temperature strength. After these measurements had been made the specimen was put back into the oil bath and the annealing treatment continued until the specimen was quenched to room temperature again to measure the new room-temperature resistivity and room-temperature strength for the larger domain size. The sequence of annealing and quenching to room temperature to measure the resistivity and strength was repeated several times until the annealing treatment did not produce any further changes in the strength or resistivity of the specimen. Jones and SykesL2 have measured the room-temperature resistivity of the disordered

OF

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CuaAu

CuaAu alloy and the value they suggest, 11.5X 1OW ohm-cm, was used to check that at the beginning of these experiments, the alloys were indeed disordered. In all cases the room-temperature resistivity of the specimens after the disordering treatment described above-remembering that prior to this treatment the specimens had been rapidly cooled from 780”C-was within the range 11.5 to 11.6X1O-6 ohm-cm. The domain size after each annealing period was estimated from the room temperature resistivity of the specimen, using the data of the domain size as a function of the resistivity published by Jones and Sykes.12 Before this could be done, however, two small corrections had to be made. In the present experiments room temperature was about 10°C higher than the temperature at which Jones and Sykes made their measurements (2O’C). The present resistivities were corrected to 20°C therefore, using the temperature coefficient of 1OV ohm-cm per “C, suggested by Jones and Sykes’ Fig. l.r? To get curves relating the room-temperature resistivity to the domain size for specimens ordered at 37j°C, 350°C and 325°C it was necessary to interpolate from Jones and Sykes’ data, which are given for specimens ordered at 376°C 346°C and 298°C. It was noticed that for large domain sizes the room-temperature resistivity is proportional to the long-range order parameter, and this criterion was adopted to get the resistivity versus domain-size curves at 375°C 350°C and 325°C (see Fig. 3*). The values of the domain sizes used in Figs. 2 and 6 were estimated from the measured room-temperature resistivities using Fig. 3. The critical resolved shear strength of the specimens was measured by straining them in tension on a very hard cantilever beam machine of the Polanyi type.13 The load and elongation were automatically recorded and an example of the type of curve obtained is reproduced in Fig. lb. In these tests care was taken to give the specimens as little deformation as possible in order 12 ,,.I

,,,,

,,,,

,,,,

,,

FIG. 3. Resistivity z~wsus domain size for annealing temperatures of 375°C 350°C and 325°C. Calculated from the data of Jones and Sykes.r2 * In the analysis of their data, Jones and Sykes say that the domain sizes they have plotted in their curves are really twice the actual domain sizes. Therefore, in Fig. 3 this factor of 2 has been included.

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to avoid any ill effects on subsequent tests due to strain hardening. In general the amount of strain incurred in each test was 0.1 per cent, which is only l/25 the extent of the jerky flow region (see Fig. la). Returning to Fig. 2, it will be seen that at first the room-temperature strength increases with increasing domain size until it reaches a maximum at around 15-30 A, then decreases until it reaches an approximately equilibrium value when the domain growth has virtually ended. In this condition the strength of the alloy is less than that of freshly quenched alloys, this observation perhaps being the same as that which led some earlier workers to conclude that ordered alloys are softer than disordered ones. When the strength reaches its equilibrium value one can regard the specimen as being comprised of very large anti-phase domains the interiors of which have the equilibrium degree of long-range order. If the specimen is now heated to a temperature lower than that at which it had previously been equilibrated, two things can happen. There may be some more domain growth, but if the specimen had been well equilibrated at the higher temperature this growth will be very slow and should be insignificant as far as the strength of the specimen is concerned. More importantly, the degree of order within the domains will increase and any change in the strength brought about by the latest heat treatment can be attributed to this change in the long range order. To examine this effect some specimens were first equilibrated with respect to the domain growth and the long-range order parameter by holding them at 375°C for a period of seven days. Previous experiment had shown that there was no measurable change in the strength or the resistivity after one day at this temperature. At the end of seven days the specimens were quenched to room temperature and their strengths measured. Subsequently they were equilibrated at a lower temperature and again quenched to room temperature to measure their their strength. Equilibration was repeated at successively lower temperatures, ranging from 375°C down to 230°C the specimens being quenched from each temperature to room temperature in order to measure their room-temperature strength. Below 230°C the ordering reaction proceeds very slowly and one cannot be sure that the specimens ever reach equilibrium. The influence of long-range order upon strength is shown in Fig. 4, which gives a plot of the critical resolved shear strength (u) versus the long-range order parameter (S) and it will be seen that as S increases the room-temperature strength decreases. The range over which S can be varied in the CuaAu alloy is very small, i.e., 0.8 to 1.0, and it is difficult to determine any quantitative functional relationship between S and u. For this purpose it would be desirable to study the relationship between long-range order and strength on an AB type alloy where S can be varied all the way from unity to zero, as was done by Green and Brown14 with p-brass.

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FIG. 4. The variation of the critical resolved shear stress with the long-range-order parameter (S). The results from four different specimens are shown in this figure. The values of S were taken from the curve S versus temperature published by Keating and WarrenF7

They showed that the strength increased as S decreased from unity to 0.5, and then decreased again as S proceeded from 0.5 to zero. Qualitatively, it is clear why the strength should rise as the long-range order parameter decreases from unity, both in CuaAu and in P-brass, and why the strength should fall again as the long-range order parameter approaches zero in &brass. Consider a long-range ordered crystal made up of a single domain in which the order is perfect. i.e., S= 1. If one passes a dislocation appropriate to the ordered crystal of the crystal lattice through this crystal, there will be no change in the atomic arrangement across the slip plane and only a vanishingly small stress will be required.15J6J7 If, however, there is not quite perfect longrange order in this domain so that some atoms are out of place, the atomic arrangement across the slip plane will be changed after the dislocations have passed through. Energy will be required to do this and the applied stress will have to be raised accordingly. The further S departs from unity the worse the misarrangement caused by slip will be, and the applied stress will have to be increased proportionately. In effect this is just Fisher-Y6 idea for short-range order hardening applied to the case of the departure from perfect longrange order rather than to the case of the departure from a complete random arrangement. Let us consider the disordered alloy for a moment. Fisheri has pointed out that when dislocations are passed through a completely disordered alloy the random matching of atoms across the slip plane is re-created behind them. The internal energy of the crystal is not changed, therefore, and consequently the stress required to move the dislocations is small. If there is some short-range order present, how-

ARDLEY:

STRENGTH

ever, it will strengthen the alloy because as the dislocations pass through the crystal the degree of short-range order across the slip plane is reduced to a more nearly random arragement of higher energy. Thus, if the arrangement of the atoms in an ordering alloy is changed continuously from a random distribution to the completely ordered distribution we would expect the strength of the alloy to vary in the following way: at first, the strength will increase as the shortrange order parameter increases, and then later it will decrease as the fully ordered state is approached. To evaluate the strength due to the domain size alone, the contribution from the long-range ord,er parameter has to be subtracted. The contribution to the strength from the long-range order alone cannot be determined experimentally over the whole range of S from zero to unity because of the interference of the domain growth. Therefore, it was decided to follow the strength as S changed over a limited range and then to assume that the strength would relax from the freshly quenched value (A in Fig. 2) to the equilibrium value (B in Fig. 2) with the same relaxation time. The following procedure was adopted to measure the relaxation times. Specimens were first disordered by heating them to 420°C for 24 hours; then they were transferred to a furnace at 375°C and held there for seven days in order to equilibrate them with respect to both domain growth and longrange order parameter. The specimens were then quenched down to some lower temperature (35O”C, 325”C, 300°C or 275°C were used) and their strength followed as a function of time at that lower temperature. The specimens were always strained at room temperature, which meant that the relaxation process was interrupted several times by a quenching treatment. Each time the strength was measured the specimens were deformed a little (approximately .l per cent tensile strain) but it was assumed that this would not have any serious effects on subsequent measurements. This is a fair assumption to make because it is suspected from previous observations that in the early stages of deformation slip is confined to discrete locations and providing we stay within the so-called “jerky range,” we will be measuring the strength of undeformed material. It was found that the relaxation of the strength of the crystals could be related to the time through the simple exponential equation :

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529

CuAur

I 1.8

I 1.7

1.6

+x10’

FIG. 5. Relationship between time of the strength of a CulAu of the temperature.

IN

I

‘K-’

the logarithm single crystal

of the relaxation and the reciprocal

port their assumption-that the relaxation time for the ordering process would be related to the temperature through a single activation energy so that T could be written 7= 70 &ET The values suggested for 70 and w/R by the straight line in Fig. 5, are 10-12.4 set and 19,5OO”K, respectively; these are very close to Bragg and Williams’ calculated values of 10-12 and 19,OOO’K. Whilst the agreement here is very good, the quality of the results could undoubtedly be improved by using the electrical resistivity rather than the strength, to follow the course of the relaxation. Estimates of the relaxation times at 375”C, 350°C and 325’C from Fig. 5 are 4,14 and 4.5 seconds, respectively. These are so low that generally by the time the first measurement was taken in the strength+ersus-domainsize experiments, the long-range order parameter had essentially relaxed to equilibrium. Under these circumstances the correction to the strength due to the influence of S degenerates into merely subtracting the equilibrium value of the strength from all the other measurements; this is shown schematically in Fig. 2. This form for the correction cannot be true in the earliest stages of ordering because there the degree of longrange order and the domain growth are intimately connected and any change in one produces an appreciable change in the other. It is worth noting here that on three occasions the strength was observed to decrease a little TABLE I.

where ui is the initial strength, rf the final strength, ut the strength at the time t, and 7(T) the relaxation time for that particular temperature. The values of r obtained from this relationship were used to construct Fig. 5. The straight line in this figure was drawn to give the best fit to the experimental points by the method of least squares. Bragg and WilliamP assumed-and the experimental results of Sykes and Evans28 tend to sup-

Specimen number

Annealing temperature

Equilibrium strength cm (kg/mm*)

Y ergs/cm~

::

325°C

2.3 2.25

;i

:

375°C 350°C

3.5 3.1

115 98

5

375°C

3.6

104

I A

3.75 3.75 3.75

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FIG. 6. Room temperature strength vwsus domain size for annealing temperatures of 37S”C, 350°C and 325°C. The left-hand ordinate scale shows the absolute value of the critical resolved shear stress.

-about 5 per cent of the vaiue for freshly quenched specimens-before it began to increase. One of these occasions is recorded in the experimental data of specimen 5 in Fig. 6. The final data obtained for the strength of an ordered crystal as a function of the domain size for three different annealing temperatures 375°C 3.50% and 325°C are shown in Fig. 6. The points on these graphs represent the experimental observations whilst the full curve is a plot of the equation fl=-

y

E--d

e (-> E

3 f

made up of anti-phase domains whose average diameter is ‘e, then it can be shown when the total slip displacement is e/2, the amount of anti-phase domain boundary in the slip plane is equal to one half of the total area of the slip plane. Any further slip does not change this ratio so that there is no longer any contribution to the strength from the ordering forces. The energy per unit area of the slip plane due to wrong bonds across the slip plane is therefore r/2(1 -[ f~-~)/e]~} before slip and y/2 after slip. Thus one can write for the work done during slip E Y /c-t\ 3

(1)

where u is the resolved shear strength of the crystal, y the surface energy of a stacking fault in the ordered structure, e the anti-phase domain size, and t the thickness of the anti-phase domain boundaries. This equation was first derived by CottrelP in a slightly different form, by equating the energy of the domain boundaries created during slip to the work done by the applied stress. The model he used is illustrated schematically in Fig. 7. Figure 7a shows an ordered structure containing domain boundaries, before slip; the atoms within the boundaries are thought to be distributed at random. Figure ib shows the same crystal after it has been slipped a certain distance and it will be seen that the amount of anti-phase boundary has been increased. The energy to do this has to be supplied by the externally applied stress and obviously the more domain boundaries there are-i.e., the smaller the domain size -the larger that stress will have to be. If the crystal is

U’--=-

22 ‘(

-

(2)

7

E J

from which Eq. (1) is derived. For eat,

Eq. (2) reduces

-aXOXOX$XOX*XOXOihXOXO

(4

$ oxox&pxoxoxox~xoxox xoxox~xo xoxoxc$oxo oxoxokxoxoxox$xoxox

x0

xoxox~noxoxoxopxoxo oxoxc$pxoxoxo%~xoxox

i

DOMAIN

-

BOUNDARY

(SCHEMATIC)

FIG. 7. Schematic illustration of the effect of slip on the anti-phase domain boundaries.

ARDLEY:

STRENGTH

OF-

CUJAU

531

to o=r/e(l-3t/e) which is analagous to Cottrell’s equation. The values obtained for y and t, from the present investigation, range from 94 to 115 ergs per cm2 and from one to two atom spacings respectively; see Fig. 6 and Table I. From their X-ray and resistivity measurements, Jones and Sykes12 also suggest that the domain boundaries are about one or two atom spacings wide. One can estimate the value of y by considering the energies of the like and unlike bonds in a stacking fault of the ordered structure. Peierls19 has obtained an expression relating these energies to the critical temperature of the alloy, and using his expression one obtains for y, when S= 1, a value of approximately 75 ergs/cm?, which agrees fairly well with the experimentally determined values. EXPERIMENTAL RESULTS III A few experiments were carried out to see how the strength of the CU~AU crystals varied above the critical temperature and the results of these experiments are summarized in Figs. 8, 9, and 10. The specimens were first ordered and equilibrated at 300°C and then their strength was measured, at temperature, over the temperature range from 300°C to 68O’C. From 3OO’C to the critical temperature there is a marked increase in strength, presumably due to the decrease of the longrange-order parameter. As the specimen passes through the critical temperature there is a sharp decrease in the strength accompanying the disordering process. This decrease is drawn as a dotted line in Fig. 8 because it was not possible to determine whether it occurred discontinuously or over a finite temperature range. Above the critical temperature the strength increases again until

SHEAR

300

STRAIN

RATE

400 TEMPERATURE

4.4X IQ’SEC:

500

600

700

OF TESTWS -C

FIG. 8. The critical resolved shear stress versus temperature for a single crystal of CuaAu. The crystal was originally ordered at 300°C before these data were taken. Shear strain rate 4.4X lo-* sec6.

FIG. 9. The critical resolved shear stress for different rates of strain. Curve D was specimens.

wrsws temperature obtained from two

it reaches a maximum value at a temperature 2”’ and then it decreases again. The position of this maximum was found to be strain-rate dependent, moving to higher temperatures the higher the strain-rate, see Fig. 9. It should be pointed out here that in addition to the change in strength with temperature there are also changes in the shear moduli and for strict comparative purposes one ought to plot strength divided by shear modulus ziersus temperature rather than just strength versus temperature. However the changes in the shear moduli whilst not negligible, are not large enough to alter the curves of Figs. 8 and 9 appreciably. Since there is only continuously decreasing shortrange order as the temperature is increased above the critical temperature, one cannot attribute the peak at 2” to a similar origin as the one at the critical temperature. However, it has been well established that in certain aging materials such as iron20~21~22 and aluminum,23 there is a temperature range over which the strength increases with increasing temperature. In iron this is comonly known as the blue-brittle range, and the increase in strength is thought to be due to strain-aging occurring simultaneously with the deformation.23 It was thought, therefore, that strain-aging could also be the cause of the peak at T’, and further evidence to support this point of view arose when two other prominent features of an aging alloy, namely an inverse rate effect (i.e., strength decreases with increasing strain-rate) and a strong yield point, were both found to occur above the critical temperature (see Figs. 9 and 10). It is not surprising that one observes a yield point and strain-aging in the disordered CusAu alloy because there are three possible mechanisms from which these effects might arise. These are (a) Cottrell locking, where the solute atoms form an atmosphere around the dislocations and pin them, 24 (b) Suzuki pinning, where the solute atoms interact chemically with the stacking faults between the partial dislocations of the face-centeredcubic lattice and pin the dislocations,25 and (c)

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FIG. 10. Showing a vield ooint and strain-a&e in a sinale crystal specimen 07 C&Au. &fore each test &e specimen was rested at 565°C for the following period of time. (A) 12 hrs; (B) 30 see; (C) and (D) 1 hr; (E) 15 set; (F) 3 hrs; (G) 1 hr. Curves A to F were obtained with a shear strain rate of 3X lo-* se@ and curve C with a shear strain rate of 7.5X10-* set‘-‘.

CottreW7 development of Fisher’s idea,26 that in an alloy in which there is short-range order, a larger stress is required to start deformation than to maintain it afterwards because the first dislocations to pass across the slip plane produce more disorder than those that follow them. If the deformation is stopped and the short-range order allowed to re-establish itself, the yield point will return and the alloy can be said to have strain-aged. These three mechanisms should apply, of course, to any substitutional solid solution alloy in which short-range order exists, and from this point of view the results obtained on Cu-Zn alloys by Ardley and Cottrellli support the findings of the present investigation. One other observation which is interesting to record is that below the temperature of the second peak, T’, deformation always occurred discontinuously, whereas above that temperature deformation was smooth. A similar effect was observed in @-brass by Ardley and Cottrell,” except that in their case the maximum in the strength-versus-temperature curve, and the transition from jerky flow to smooth flow occurred below the critical temperature. CONCLUSIONS

It has been shown that the room-temperature strength of an ordered alloy of CuaAu depends on both the degree of long-range order and the anti-phase domain size. It is suggested that this dependence might be predicted by considering the change in bond energies across the slip plane which is brought about by slip. It is now clear how earlier workers might arrive at different conclusions as to the effect of ordering upon the strength of alloys. Those who had well-ordered alloys would observe a low strength whilst those who had either incomplete ordering or small domain sizes, or

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both, would observe a high strength. This is particularly important when comparing the CusAu type alloys with the CuAu types where the ordering rates are so very diBerent. The “transition range” mentioned in the early literature on CuAu alloys is presumably just the period in which noticeable domain growth is occurring. In view of the large effects the domain size and the long-range-order parameter each have on the properties of ordering alloys, one ought to be careful to separate them when investigating the disorder+order reaction. In the disordered region, yielding, strain-aging, an inverse strain-rate effect and an inverse temperature dependence of the strength upon temperature have all been observed, thereby forming a picture which is consistent with the current ideas of the effect of solute atoms and their arrangement in the lattice, upon the plastic deformation of alloys. ACKNOWLEDGMENTS

I wish to thank Dr. J. C. Fisher and Prof. A. H. Cottrell for their advice and encouragement throughout the course of this work. I would also like to thank Mr. V. J. DeCarlo for helping me to prepare the single crystals. REFERENCES L. Nowack, Z. Metallk. 22,94 (1930).

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