AMERICAN INSTITUTE OF MINISG AND METALLURGICAL EICGIICEERS Technical Publication No. 2341

Class E, Metals Technology, February 1948 DISCUSSION OF THIS PAPER IS INVITED. I t should preferably be presented by the contributor in

Person at the New York Meeting, February 1948. when an abstract of the aper will be read. If this is impossible. discussion in wntlng (a copies) may be sent to the Secretary, Amencan fnstitute of Mining and Metallurgcal Engineers, a9 West 39th Street. New York 18. N. Y. Unless special arrangement is made. discussionof this Paper will close May 15. 1948. Any discussion offered thereafter should preferably be in the form of a new paper.

Mechanism of Precipitation in Alloys of Beryllium in Copper* BY A. G. GUY,^ C. S. BARRETT,~ MEMBER, AND R. F. M E B L , ~ MEMBER AIME

(New York Meeting. February 1948)

INTRODUCTION was found. Aging and overaging were ob-

IN the last few years this laboratory has served to occur most rapidly in the region

published a series of papers on the of discontinuous precipitation-the grain nism of age-hardening.i4,i.6,~1,~o ~ ~ i ~ f l ~ boundary-apparently accelerated there by

stated it has been proposed that hardening strain.

is caused by lattice strains created by lat- The constitution the c ~ - B ~ system tice coherency between matrix and precipi- tate; much experimental evidence has been The constitution diagram is advanced to support this thesis. shown in Fig I. The structure of the y

The system c ~ - B ~ is an one to phase was determined by Misch16 as of the

study. It is a system that exhibits marked caesium-chloride type with a lattice con-

age hardening; it ought to be subject to the stant Of "598 A. High temperature

type of analysis of structural changes photograms Of 7.2 pct Be made by

hitherto explored and it is a prominent Koss~lapow and Trapesnikowl' led them

example of u~~~~~~~~~~~~~ precipitation,,, to conclude that the structure of the B which, though studied in other cases,6.2~

phase is simply disordered body-centered

remains largely unexplained. cubic with a0 = 2.79 A a t 750°C.

This research shows that hardening in Prenious Work the Cu-Be system proceeds essentially as

The initial work on the age-hardening of proposed earlier.14 Guinier-Preston zones Cu-Be alloys was that of Masing, Dahl, have been observed to form on the (1001 Holm, and Haase. Their numerous papers planes of the matrix; no transition lattice were ~ublished in book form in 1929, sub- . . .

This paper represents part of a thesis sub- sequently translated.18 The effects of Be mitted b y A. G. Guy t o the Graduate Com- mittee of the Carnegie Institute of Technology, content, SOlUtiOn temperature, aging tern- Pittsburgh, Pa.. in partial fulfillment of the re- perature, and cold work on the hardness quirements for t h e degree of Doctor of Science. May. 1946. Manuscript received a t the office of developed in these alloys were studied. An the Institute Sovember 17. 1947.

+ Associate Professor of Metallurgy, Mechan- initial decrease in the electrical conduc- ical Engineering Department. North Carolina tivity was observed when strips of a 2.5 pct State College of Agriculture and Engineering, Raleigh, North Carolina. Formerly Metals Re- Be alloy were aged at temperatures from search Laboratory Graduate Fellow, Depart- 200 to 3 5 0 0 ~ ; a wire of the same analysis ment of Metallurgical Engineering, Carnegie Inst i tute of Technology. Pittsburgh, Pa. failed to exhibit this decrease when aged a t

S Professor, Institute for t h e Study of 3 5 0 0 ~ A volume decrease of 0.6 oc- Metals, University of Chicago. Chicago. Illi- nois. Formerly Staff Member. Metals Research curred during aging a t 250°C. Laboratory, and Associate Professor. Depart- ment of Metallurgical Engineering, Carnegie Mi~rOscO~ic investigations Of quenched Inst i tute of Technology. Pittsburgh, Pa. and aged alloys disclosed a regularly ori-

O Director, Metals Research Laboratory, and Head. Department of Metallurgical Engineer- ented "striping" (ripples) on the a and on ing. Carnegie Institute of Technology. Pitts- the B grains. In the case of the P the cause of burgh. Pa.

1 4 References a re a t t h e end of t h e paper. the striping was held to be decomposition

Copyright. 1948. by the American Institute of Mining and Metallurgical Engineers, Inc. Printed m USA

2 MECHANISM OF PRECIPITATION IN ALLOYS OF BERYLLIUM IN COPPER

of the /3 grains. Incipient ~recipitation of Desch3 repeated some of these micro- y crystals in the a solid solution was sug- scopic studies and obtained the same pe- gested as the cause of the striping in the a culiar markings on the aged specimens. He grains. The grain striping was found to concluded that submicroscopic particles

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reach its maximum intensity after a period were precipitating on the octahedral slip of heating corresponding to maximum planes of the Cu-Be alloy. hardness. B ~ m m , ~ using an alcoholic iron-chloride

X ray studies were confined to the use of etching solution, was able to show a dark- a Debye camera on p~l~crystalline wire specimens 0.5 mm in diam. From the data obtained it was concluded that the equilib- rium precipitate CuBe appears before the hardness maximum has been attained. The sharpness of the matrix lines was reported to decrease during the early stages of aging, increasing as the specimens were overapd. Two possible causes of this line widening were suggested, namely, variation in solute concentration of the matrix, and internal strain. From the results of experiments it was decided that lattice strain is the pri-

etching grain boundary material. At lower aging temperatures this grain boundary material was observed to grow a certain distance into the grain and then to stop. On aging at 5o0°C, however, it grew to cover the whole grain. The halting of the grain boundary reaction was interpreted as indi- cating the presence of a submicroscopic reaction in the center of the grain.

Swindells and SykeslB used a 2.6 pct Be alloy to determine curves of specific heat vs. temperature for slowly cooled and for quenched specimens. On the basis that the

mary cause of this widening. energy associated with lattice distortion in

A. G. GUY, C. S. BARRETT AND R. I?. MEHL-TP 2341 3

severely cold-worked metals is about 0.5 that they represent the intersection of cal per g, and that the chemical energy (IIO] planes with the polished face. The available in the supersaturated solution is appearance of these streaks was attributed about 6 cal per g, any mechanism that to a mechanical action associated with the allows as little as 10 pct of this chemical precipitation mechanism. energy to be stored as lattice strain will be From the results of X ray data obtained effective in hardening the metal. From the on single crystals using monochromatic fact that maximum hardness is attained in radiation, the following precipitation mech- the Cu-Be alloy after the release of the anism was proposed: major portion of the chemical energy, these I. The assembling of Be atoms on (IOO] workers concluded that precipitation is re- planes of the matrix to form Guinier- sponsible for the hardening observed in this Preston zones. system. I n contrast to this Jones and 2. The progressive formation of the y Leech12 interpreted their specific-heat ex- phase, CuBe. Marked increase in hardness periments as indicating that some form of is caused by the lattice changes associated transformation occurs in addition to the with this stage. main transformation associated with the 3. The third stage is the appearance of precipitation process. Jones and Leech also the y phase precipitate in the a matrix and concluded from the results of Debye X ray is accompanied by a decrease in hardness. experiments that the equilibrium precipi- The orientation relationship is nearly tate is present in alloys showing maximum hardness.

After the present work had been com- pleted two by Guinier and J a c q ~ e t ? . ~ became available in this country. A brief summary of their work is given a t this point and reference will be made to their results in several later sections of this paper.

The alloy studied by Guinier and Jacquet was of commercial quality, containing 2.3 pct Be, 0.2 pct Fe, 0.1 pct Si, and 0.2 pct Mg. That the /3 phase present in the quenched specimens of this material had no appreciable effect on the reaction mecha- nism was shown by companion experiments

Guinier and Jacquet concluded that con- trary to the views of previous workers the Cu-Be system is not an example of a system that shows equilibrium precipitate suffi- ciently early in the aging process that har- dening can be attributed to its action. Instead they class it with A1-Cu and A1-Ag as regards precipitation mechanism.

Preparation of Alloys

on a 1.9 pct Be alloy. Single crystals pre- The polycrystalline alloy was melted in a pared by slow cooling from the melt and graphite crucible in a gas-fired furnace polycrystalline sheet one mm thick were under a sodium chloride flux. The analyses used in obtaining the following results. of the materials used were:

I I I

Master alloy.. ............................ 4 . ar pct Be 0 . 0 6 pct Pe bal. Cu OPHC Copper.. . . . . . . . . . . . . . . . . . . . . . . . . . . o .ooa pct Ni 0 .003 pct Pe 0 .02 pct 0 2 99.97 pct CU

Micrographs of electropolished aged The melt was stirred with wood and with specimens showed grain boundary precipi- graphite rods and was cast into a vertical t i te and also streaks in both a a n d p grains. iron mold. Contrary to the observation by Desch that Alternate cold-rolling and annealing were the streaks in the a grains are (111) plane used to reduce the thickness of the ingot traces, Guinier and Jacquet determined from 1% to 36 in. This material was used

4 J IECIUSISJ I OF PRECIPITATION IS ALLOYS OF BERYLLIUM I N COPPER

for hardness and electrical conductivity specimens. The wire used for obtaining nebye S ray patterns was drawn to 0.028 in. The grain size of this \\ire was about the same as that of the In-in. stock.

Single crystal specimens were desired for Laue S ray patterns and these were made by the Bridgman technique. Two satis- factory specimens were made using the melting procedure outlined above.

The chemical analyses of the three allays \\ere as follows:

Per Cent - - .~~ -

I Rc ; Fe Si Cu - - - .

l i l ! Pr,lycrystallinc all,,y.. . r 92 0. 02 o 01 b.11. Sinrrlc crystal.. . . . . I . 73 o 07 o. 0 1 h 3 1

Tri-crystal.. . . . . . . . . . ; I 59 , 0.04 I 0.01 8 bll.

have been found for A1-Cu, A1-Ag, Al-Zn, and A1-JIg alloys. -4 search for such streaks was made on the Cu-Be single crystal.

Thin disks from the single crystal were cut into pieces about 34 by in. and were further thinned by alternate polishing and etching do\vn to 0.005 in. After being sealed under vacuum in individual quartz tubes the specimens were heated a t 8o0°C for 116 hr and were quenched directly to the aging temperature. Since all of the heating was done by means of salt baths the heat- treating temperatures were quickly reached and the quenching was moderately fast. I t has been shown that the quenching of Cu-Be alloys need not be especially rapid for satisfactory retention of the solid solu- tion.18 Air-cooling was used from the aging temperature.

IVhen it was time for a sample to be used the quartz tube was broken and the speci-

:I transition lattice has been shown to men was mounted in the Laue ray cam- play an im~)ar tant part in the mechanism era. An iron filter one mil thick was placed of precipitation hardening in the .\l-Cu and immediately in front of the Eastman in the A1-Ag systems. Since previous double-coated So-screen film used for this workers had failed to find such a lattice in work. The Coolidge tungsten ray tube Cu-Be alloys using the usual Ilebye S ray ,\-as at 35.000 and milli- technique, improved methods were used in amI)eres. To produce films I\.ith strong the present investigation. The refinements Guinier-I,reston streaks an exposure time employed were a monochromated S ray of 2 days was used. beam, an improved system of tlafies, and The Laue pattern obtained from a speci- evacuation of the air from the S ray men aged for Iooo hr at 2000C in camera. Although a series of films corre- Fig 2. Fig is the stereographic projection sponding to various aging times was made of some of the spots and streaks of Fig z . at each of the temperatures 3O0! and I t is seen that the Guinier-Preston streaks 4 0 ~ " ~ trace a lattice was follow great circles through the I loo 1 poles. observed. The equilibrium precipitate was From the analysis developed by Barrett

in aged the longer and Geislerl it is apparent that this indi- times at 30° and 4 0 " ~ as rel)Orted cates the formation of two-dimensional investigators (see above). plates parallel to the 11oo) planes of the

matrix. Therefore it can be concluded that Larre S ray Stzrdies

Guinier-Preston zones form on the { r o o ) Strong evidence concerning the mecha- planes of the matrix during the course of

nism of precipitation hardening is given by precipitation in Cu-Be alloys. the presence of certain streaks on Laue I n order to study the development of S ray patterns. These streaks are indica- the Guinier-Preston zones with increasing tive of Guinier-Preston zones and they aging time a series of Laue photographs

A. G. GUY, C. S. BARRETT AND R. F. MEHL-TP 2341 5

were obtained from specimens aged for fracted spots is in marked contrast to the various times a t 2 0 0 , 300, and 4o0°C. As appearance of those obtained from the aged reference patterns Laue photographs were specimens. also made using a Cu single crystal, Fig 4, Fig 6 shows the Laue patterns obtained

FIG 2-GUINIER-PRESTON STREAKS IN THE FIG 4-LAUE PATTERN PROM A CU SINGLE

L A U E PATTERN OP A SPECIMEN AGED 1,000 HR CRYSTAL. AT 200°C.

FIG 3-STEREOGRAPHIC PROJECTION O F SOME O F THE SPOTS AND STREAKS O P F I G 2 .

and a water quenched Cu-Be specimen, Fig 5. NO curved or center streaks1 indicac- tive of Guinier-Preston zones were visible in Fig 4 although thermal streaks are in evidence. In Fig 5, too, thermal streaks are present with perhaps a slight indication of center streaks. The sharpness of the dif-

FIG 5-LAUE PATTERN FROX A AVATEX QUENCHED CU-BE SPECIMEN.

from samples aged x, 19, 200, and 1,000 hr respectively a t 2o0°C. The orientations of the crystals were such that the Guinier- Preston streaks are predominantly radial, but by using the water quenched specimen and the half-hour specimen as references the importance of radial thermal markings

If'

I - 1

FIG 6-LAUE PATTERNS FRO11 SPECIMENS AGED FOR VARIOUS LENGTHS OF TIME AT 2 0 0 ° C . a. 0 .5 hr, b. 19 hr. c. 2 0 0 hr. d. 1,000 hr.

TIME /N HOURS FIG 7-HARDNESS, ELECTRICAL RESISTANCE, AND PER CENT GRAIN BOUNDARY REACTION CURVES

FOR POLYCRYSTALLINE ALLOY AGED AT 2 0 0 ° C .

A. G. G U Y , C. 8. B,\RRETT AND R. F . MEHL-TP 2341 7

can be assessed and proper allo\vance made. hours. Xo trace of the equilibrium precipi- The streaks going towards the center of the tate is present even after aging for 1,000 hr. film are prominent but appear to change In Laue patterns from specimens aged a t little for aging times greater than 19 hr. jm°C gradual development of the Guinier-

TIME IN mURS FIG S -IFARDXESS, F;LECTRIC,ZL RYSISTASCE, , n n PER CEST GRAIS B O ~ ~ A R Y REACTIOS CCRVES FOR

POLYCBYSTALLISE ALLOY AGED AT 300'c.

TIME /N W R S FIG ~- I~ARDSKSS A S D ELECTRICAL RESISTANCE C L W E S FOR POLYCRYSTALLISE ALLOY A(;ED h T

350°C.

This series of patterns indicates that Preston zones could again be followed. Guinier-Preston streaks appear after a few Diffraction spots from three-dimensional hours aging a t 2o0°C and continue to precipitate were observed in the pattern increase in intensity with continued aging from the roo hr specimen. On aging a t reaching a maximum after a few hundred 400°C diffraction effects from three-dimen-

sional precipitate were present after 4 hr solution treatment. Resistance readings on though the Guinier-Preston streaks had not the specimen and a standard resistance disappeared after 26 j hr. were sufficient to determine a "compara-

tive resistance" in arbitrary units. ELECTRICAL KESISTASCE JIEASL*REMESIS ~ h , results obtained for the polycrystal-

The specimens used to determine the line specimens are 111otted in Fig 7-10. resistance change on aging were >8-in. sq The initial anomalous increase in resistance

TIME /N mlURS FIG 10-~I.IRDSESS, ELECTRICAL RESIST.ISCE, ASD PER CEST GRAIS BOCSDARY REACTIOS CCRL'ES

FOR POLYCRYSThLLISE ALLOY AGED AT 400°C. 7- I

T/ME /N HOURS -H.\RDSESS ASD ELECTRICAL RESISTASCE CI.B\-ES FOR SISGLE CRYSTAL AGED

polycrystalline rods about 3-in, long. These is seen to decrease in magnitude a t higher \rere quenched directly in the aging salt aging temperatures, as found b y earlier bath from a salt bath at 800°C used for the workers, vanishing a t 350°C. The addi-

A. G. GUY, C. S. B.\RRETT AND R. F. AIEHL-TP 2341 9

tional curves in these figures r i l l be dis- tures, this efiect must be caused by some cussed in subsequent sections of this paper. other factor, one that is accelerating the

Resistance measurements irere also made aging rate. .\ series of tests showed that this on ? 6 by ?4 in. disks cut from the single other factor is prot~ably the stresses pro-

5Wt , . . , . , . I , . , . .&LL A II I A. L A20 0.01 0.1 I I0 100 IGm

TIME /N HOURS FIG I 2-HARDSESS AND ELECTRICAL RESISTASCE CTRVES FOR SISGLE CRYSTAL .AGED .AT 300°C.

crystal. The method of testing was the same as that used for the polycrystalline specimens. From the results shown in Fi,y I 1-13 it is seen that the aging reaction proceeds less rapidly than in the polycrys- talline specimens as has been observed in other alloy systems.

Hardness Jfeasureiiteitts

Specimens for hardness measurements were cut from the polycrystalline ?$-in. square rod and were about 94-in. long. They were given the usual 1>4 hr solution treat- ment a t 800°C followed by direct quenching to the aging temperature. Air-cooling was used from the aging bath. The average of six Rockwell A scale readings are shown in Fig 7-10.

One peculiarity noticed in these hardness curves is that the curve for zoo0C begins a t a markedly greater value than do those for higher aging temperatures. Since the reaction velocity is lower at low tempera-

duced in the specimen as a result of quenching to the aging temperature. These stresses arc presumably greater the lower the quenching (aging) temperature and thus tend to accelerate the aging of speci- mens quenched to lower temperatures.

Fig 14 gives a comparison of the hardness of a specimen water-quenched to room temperature and of one hot-quenched to the aging temperature. Although the water-quenched sample begins its reaction at a higher rate the hardness value it reaches is only one point greater than that attained by the other.

To measure the change in hardness of single crystal specimens, disks from the single crystal were given a 1%-hr solution treatment at 8o0°C follo\red by quenching to the aging temperature. At each tem- perature only one disk was used for the entire series of measurements, the orienta- tion of the specimen in the hardness testing machine being maintained constant to

I 0 M E C W S I S M OF PRECIPITATION I N ALLOYS OF BERYLLIUhf I N COPPER

avoid orientation eEects. Before testing, An interesting phenomenon was observed the surface was given a metallographic on the single crystal sample aged a t 3o0°C polish through 600 carborundum on a for times longer than 3 h r s t r e a k s oc- canvas wheel. Fig 11-13 present the curred on the surface of the polished

5 d O l ' ' - 0.1 ' " " " ' i ' ".- I 0 ' " ' - 100 . TIME IN HOURS

FIG 13-HABDSESS ASIJ ELECTRICAL RESISTASCE CURVES FOR SISGLE CRYSTAL AGED AT 4oo°C.

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1

TIME /N HOURS FIG I~-COSIPARISON OF HARDNESS A h n PER CEXT GRAIN BOVXVARY REACTIOS CURVES FOR WATER

Q n S C H E D ASD HOT QL7ESCHED SPECDLESS ACED AT 300'c.

results obtained using the Rockwell Super- specimen on being removed from the aging ficial 1 5 s scale. .\ comparison of these bath. These streaks must have been caused curves with those obtained for the poly- by displaced metal for they disappeared crystalline specimens shows that the single when the specimen was repolished. I t was crystals harden much more slo\vly a t all possible to see three streak directions with temperatures. the unaided eye though attempts to

A. G . GUY, C. S. BARRETT AND R. F. MEHL-TP 2341 11

observe the streaks under the microscope reflection Laue photographs. By suit- were unsuccessful. able heat-treatment the "displaced-metal

Because of the importance of these streaks" were caused to appear on the streaks with respect to the mechanism of polished surfaces of these disks and the

3 0 L

C /. I ,------A\: '.

0 --- GRAIN I GRAIN n . - 6 R A I N l U B w m w - G R A I N m

om 0.1 10 l a , lo00

TIME IN HOURS FIG 15-MICROHARDNESS OF THE GRAIN CENTERS AND OF THE GRAIN BOUNDARY AREA OF GRAIN

111 I N THE TRICRYSTAL AGED AT 300°C.

precipitation, two additional experiments streak directions were compared with were completed in an effort to determine the directions of the traces of [ 100) , ( IIO 1, the crystallographic plane on which the and (1111 planes. I t was concluded that metal was being displaced to form the the streak directions did not coincide with streaks. The orientations of two single the traces of any one of these sets of crystal disks were determined using back- planes. This result is at variance with that

150

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TIME IN HOURS FIG I~MICROHARDNESS OF GRAIN 11 I N THE TRICBYSTAL AGED AT 400°C.

I 2 MECHANISM OF PRECIPITATION I N ALLOYS OF BERYLLIUM I N COPPER

obtained by Guinier and Jacquet who also studied this effect; it also indicates that the streaks are not simply the traces of slip planes, i.e., ( I I I J planes in face-centered cubic metals. However the result is not inconsistent with the fact discussed below, that the orientation relationship of the precipitate and matrix is complex rather than simple.

Microhardness

In order to determine the concurrent change in hardness of the grain boundary and of the grain center areas during the course of aging, it was necessary to use a tricrystal so that crystallographic orienta- tion could be eliminated as a variable. Disks from the tricrystal, ground so that the opposite faces were accurately parallel, were water quenched after being heated for 1% hr at 800°C. After careful mechani- cal polishing the disks were tested on a Tukon machine using a 2 0 0 g load. .

The changes in microhardness of the grain center and grain boundary areas of a single specimen on aging a t 300 and 4m°C are shown in Fig 15 and 16. At the lower aging temperature the grain boundary region apparently was overaged by the time it was large enough to be measured. At 400°C however the first measurable grain boundary area was considerably harder than the grain-center. With con- tinued aging the grain-boundary material became as soft as or softer than the balance of the crystal. A reasonable explanation of these results seems to be that faster aging and faster overaging occur a t the grain boundaries.

During the course of these experiments on the microhardness of the tricrystal a number of interesting observations were made. Fig 17 is a micrograph of the specimen aged 31 hr at 200°C. The peculiar markings that look somewhat like small recrystallized grains appeared on etching specimens that had been aged a short time (I hr at 4o0°C) but failed to appear on

specimens aged a longer time (longer than 10 hr a t 400°C). Since maximum hardness is attained after about 10 hr a t 400°C this phenomenon may be the evidence for reaction occurring in small domains within each crystal. In Fig 1 7 the hardness of the area inside the "recrystallized grains" was greater than that of material making up the broad "grain boundaries."

Although the 400" specimen completed its aging without accident, both the 200

and 300°C specimens fractured along the grain boundaries. In the 200°C sample the first crack was observed after 30 hr; in the 300°C sample after 8 hr. Fig 18 shows a variety of cracks in the 300°C specimen aged for 8 hr.

Two interpretations of this phenomenon are possible. One is that the reaction at the grain boundary generates stresses that cause the fracture to occur. This explana- tion is embarrassed by the fact that a great deal of grain-boundary precipitate appears on aging a t 400' without accompanying fracture. Moreover, cracking was observed at some grain boundaries in the absence of appreciable precipitation. A more probable interpretation is that the stress is generated by reaction occurring throughout the grain, the grain-boundary representing an area a t which stresses are intensified. These heightened stresses a t the grain boundaries account for the appearance of precipitate in this region and for the generation of cracks. At high temperatures the stresses are not so great as shown by the lower maximum hardnesses attained. This ac- counts for the absence of cracking at 400". That stresses account for the progress of grain-boundary precipitation is indicated by the fact that once cracking has occurred no further growth of the grain boundary precipitate is observed. A further support for this view is given by the observation that precipitation does not occur a t the boundary between two grains having similar orientations as judged by etching characteristics. The recent work of Forsyth4

A. G . GUY, C. S. BARRETT AND R. F. MEHL-TP 2341 I3

FIG 17 ( ~ ~ O V ~ ) - M A R K I N G S OBSERX'ED INSIDE A SINGLE GRAIN ACED FOR 31 HR AT 200°C. I 0 0 X. FOUR TUKON INDENTS ARE SHOMJN.

FIG 18 ( ~ ~ ~ o w ) - G R A I N BOUNDARY CRACKS I N THE SPECINEN AGED AT 300°C FOR 8 HR. 2 5 0 X. Etchant-Acid d i c h r o m a t c i solution

14 MECHANISM OF PRECIPITATION IN ALLOYS OF BERYLLIUM IN COPPER

on grain boundary precipitation in Cu-Be and Mg-AI alloys suggests that this dependence of precipitation on the relative orientations of the two grains separated by the grain boundary is a general phenomenon.

Metallog~aphic Studies Metallographic specimens were obtained

by cutting small pieces from the hardness specimens. These were mechanically pol- ished using the technique developed by Harker and Murphy."JA

Fig 19 shows the structures developed on aging at 2m°C. Fig 19a shows the phase stringers that were present in a few of the specimens despite the care taken to insure thorough homogenization. I t is seen that in the grain areas that have not participated in the grain boundary reaction there is no evidence of precipitation. A possible excep- tion may be the specimen aged for 3,000 hr in which there is a suggestion of the "ripples" discussed below. Fig 7 gives a comparison of the hardness and the per cent grain boundary reaction (the per cent of the area of the polished face that is occupied by the structure that forms in the vicinity of the grain boundary).

From the micrographs in Fig 19 for alloys aged a t 2o0°C it may be seen that the width of the dark bands at the grain boundaries is not uniform. The explanation advanced in the section on microhardness of single crystals might be applied here, namely, that the combined effect of the orientations of the two grains on either side of the boundary determines the grain boundary stress produced and hence the amount of grain boundary reaction. The peculiar distribution of the grain boundary precipitate on the periphery of many of the grains lends a measure of support to this explanation.

Micrographs of some of the 300°C hot- quenched specimens are shown in Fig 20.

While the growth of the grain boundary area is similar to that shown by the 2m°C series, there is a difference in the structure

of the grain center. "Ripples" very similar to those observed by Harkerlo in Au-Cu alloys begin to appear as the alloy ap- proaches maximum hardness. These ripples increase in prominence as the alloy overages until finally the grain-center area in which they are visible is engulfed by the ad- vancing grain-boundary region. The hard- ness and percentage grain boundary reaction curves are shown in Fig 8.

At 4o0°C the microscopic evidence of aging is of a different character as shown in Fig 21. While no ripples are seen, traces of precipitate appear in the grain center before the grain-boundary area has covered the entire sample. Fig 10 includes the hardness and percentage grain boundary reaction curves a t 4o0°C.

The pearlite-like structure that forms a t the grain boundaries after long aging a t this temperature, Fig 21d, is similar to the structure found in A1-Ag and other alloys. Judging from the apparent manner of growth of the patches, there seems to be a solution and redeposition of the precipitate a t the advancing interface as pointed out by Geisler,6 however improbable this might appear. In the A1-Ag system the transition lattice is the "dissolving" phase and the equilibrium lattice is deposited. Many bits of evidence indicate that this is not so in Cu-Be alloys. A transition lattice has not been detected; the equilibrium precipitate lines appear in the absence of this pearlitic structure a t 3m°C; and much softening has occurred before the appearance 'of this structure a t 4o0°C, showing that the equi- librium precipitate has formed in appreci- able amount. Micrographs have been obtained showing what appears to be the encroachment of the large grain boundary precipitate particles on the smaller needles in the grain center. I t would seem that in this system the equilibrium precipitate dis- solves a t the advancing "pearlitic" inter- face and redeposits on the larger plates. A possible driving force for this reaction might be the high stresses, and therefore

A. G. GUY, C. S. BARRETT AND R. F. MEHL-TP 2341 I5

FIG 19-MICROGR~PIIS OF ALLOYS AGED FOR VARIOUS TINES AT zoo°C AFTER HOT-QUENCHING TO zoo°C PROM 8o0°C. 200 X.

The corresponding reaction curves are shown in Fig 7. a. 11 hr b. 30 hr c. 198 hr d. 500 hr e . 1,000 hr f . 3,000 hr

higher energy, in other parts of the grain r = a UP (- &) Eq r compared with those in the areas that have

undergone this reaction. where Precipitation a t 560°C occurs only in a, Q = constants

small part a t the grain boundaries and R = gas constant mostly in the grain-centers. The grain T = absolute temperature

FIG 20-~'~ICROGRAPHS OF ALLOYS AGED FOR VARIOUS TIVES AT 300°C AFTER HOT-QUENCHING T O 3o0°C FROM 800°C. 2 0 0 X.

The corresponding reaction curves are shown in Fig 8. a. 0.33 hr b. 3 hr c. 7 hr d. 42.5 hr

boundary reaction is arrested when only 20

pct completed in contrast t o the IOO pct grain boundary reaction that is observed a t Iower aging temperatures.

The rate of a large variety of reactions can be expressed as a function of tempera- ture by means of the equation

If as a measure of the rate of reaction the reciprocal of the time for half-reaction is chosen, and if this quantity is -plotted against the reciprocal of the absolute tem- perature, the activation energy Q can be obtained from the slope of the resulting straight line. I n the instance of precipita- tion reactions Q is not a single quantity, but rather a complex of a number of separate activation energies."

A. G. GUY, C. S . BARRETT AND R. F. MEHL-TP 2341 I 7

Data for determining values of the ac- also holds in the present instance, taking tivation energies are listed in Table I and a the Q for electrical conductivity in the representative plot of these data is shown polycrystalline alloy as the significant in Fig 22. The values of the activation value.

FIG 21-MICROGRAPHS OF ALLOYS AGED FOR VARIOUS TIMES AT 4o0°C AFTER ROT-QUENCHING TO 4oo°C FROM 8o0°C. 200 X.

The corresponding reaction curves are shown in Fig 10. a. 0.01 hr b. 0.1 hr c. 0 . 2 hr d. 1,000 hr

energies obtained from the data of Table I ATOMIC-CRYSTALLOGRAPHIC MECHANISM OF - are given in Table 2. PRECXPITATION

Since the mechanism of precipitation I t is not possible to give a complete and hardening involves diffusion another ac- wholly satisfactory treatment of the pre- tivation energy that is of interest is that for cipitation mechanism in this system be- the diffusion of beryllium in copper. Rhines cause of our inability to demonstrate the and Mehl17 have determined this Q value occurrence and to determine the structure to be about 28,000 cal. Jetter and Mehl" of a transition lattice, as well as the inher- have pointed out that in every instance of ent complexity in the crystallographic rela- precipitation investigated the Q value de- tionships. For these reasons the folldwing termined has had a lower value than the mechanism must be considered tentative corresponding Q for diffusion. This relation; only.

TABLE I-Data for Determination of Activation Energy (a

Data from Electrical Resistance Measurements on Polycrystalline Specimens T---

Data from Electrical Resistance Measurements on Single Crystals

L ~ I I L ~ .

OC

250 300 350 400

Data from Hardness Measurements on Single Crystals

lemp. OC

300 400 560

Data for Grain Boundary Reaction in Polycrystalline Specimens

%

750 762 765 862

T

523 573 623 673

Temp, OC

300 400 560

Initial Resistance

1,000 1,000 1.000 1,070

I - T

0.001912 0.001745 0.001605 o. 001486

T

573 673 833

alloys.

15.000

Final Resistance

5 00 525 530 65 5

Time, Hr

43.7 3.75 I . 14 0.20

Initial Resistance

0.001745 1.005 0.001486 1,010 O.OOI~OO 1,075

Temp. 'C

aoo 300 400

TABLE a-Values of Activation Enevgies ferent from that of the matrix. These Given by Data of Table I platelets or "zones" give rise to the

Guinier-Preston streaks in X ray flms. Although it has been impossible to deter- mine experimentally the structure of these Guinier-Preston zones it is reasonable to assume that they form on the cube planes of the matrix from concentrations of Be

- atoms. A possible structure for the Guinier-

The first step in the precipitation reac- Preston platelets is shown in Fig 24 in tion is the formation in the homogeneous relation to the matrix, a! phase, and to the matrix of platelets having a structure dif- equilibrium precipitate, y phase.

ke 0.0229 0. I 267 0.877 5.0

Time' Hr Time

16.0 o. 0625 0.1526

0.10

Final Resistance

560 630 925

T i

?6 "lue

780 820

1.000

Time (for Maximum Hardness). Hr

40 4.7 0.65

573 673 833

T

pp

473 573 673

- Time I --

0.025 0.213 1.54

O.C"JI745 0.001486 O.OOI~OO

I - T

o.ooa114 0.001745 0.001486

Time (for 50 Pct Transformation). Hr

a.ooo I0 0.10

I - Time

0.0005 0. I0 5.0

4. G. GUY, C. S. BARRETT AND R. F. MEHL-TP 2341

0.00220

om210

000200

om90

-1: -1-

QQO(m

O.OOI50

mots OmlmWo o.ooo1 a001 M I TIYE I I YRS) 0.1

FIG 22-PLOT TO DETERMINE THE ACTNATION ENERGY FROM DATA ON GRAIN BOUNDARY REACTION-

I

1100)

FIG 23-STEREOGRAPHIC PROJECTION SHOWING APPROXIMATE POSITIONS OF CUBE (100) POLES IN THE STANDARD PROJECTION OF A MATRIX CRYSTAL.^ ( 100) POLES OF THE MATRIX ARE INDICATED BY THE POINTS WITH INDICES ATTACHED.

2 0 MECHANISM OF PRECIPITATION I N ALLOYS OF BERYLLIUM IN COPPER

Guinier and Jacquet concluded from a will result in a contraction of the lattice in study of the Guinier-Preston streaks that the direction perpendicular to the planes this intermediate structure does not exist of beryllium atoms and in an expansion in in any appreciable amount. However a the plane of these atoms. Accommodating

POSS18LE INTERMEDIATE STRUCTURE

0 COPPER ATOM

BERYUIUM ATOM

FIG 24-RELATION OF POSSIBLE INTERMEDIATE STRUCTURE TO THE a AND 7 PHASES.

comparison of the X ray effects found dur- ing the early stages of aging in Al-Ag alloy^,^ a system in which a transition lat- tice is eventually formed, with the X ray effects found in the early stages of aging in Cu-Be, reveals no significant difference. Therefore it seems likely that such an intermediate lattice may form in small amounts even though it fail to grow large enough to produce three-dimensional dif- fraction effects.

Evidently this intermediate structure will have a strong tendency to assume the dimensions of the y phase. This tendency

strains will occur in the matrix lattice and hardening will result. This hardening can occur only so long as the intermediate structure and the matrix are coherent for only then can appreciable reciprocal strain- ing exist. When the equilibrium precipitate forms from the intermediate structure, lattice strains are materially reduced and the hardness decreases.

If the equilibrium precipitate were to form from the intermediate structure shown in Fig 24 without a change in ori- entation, as observed in other systems and presumably general, the resulting rela-

A. G . GUY, C. S. BARRETT AND R. F. MEHL-TP 2341 2 1

tionship to the matrix would be described by:

This is the orientation reported by Guiner and Jacquet .s

The extent to which this relationship b differs from the one given in Fig 23 can be

-appreciated by considering the following manipulations of the stereographic projec- tion that result in approximately the pat- tern shown in this figure. Beginning with the standard projection of the precipitate coinciding with the standard projection of the matrix, the former is rotated 45" about its [oor] axis. This produces the relationship given above. Next the [oor] axis is tilted 8" about the [roo] axis. Finally by a rotation about the [oro] axis the [oor] axis is made to spread over an arc of about + 2 degrees. Efforts to explain in terms of crystal- lographic relationships why the equilibrium precipitate assumes this orientation have not been successful.

The mechanism of precipitation harden- ing proposed earlier6 proposes that: (I) hardening be concurrent with the appear- ance of a new lattice coherent with the matrix which, when in small particles, provides Guinier-Preston zones, and (2)

that softening be concurrent with the disappearance of coherency between the matrix and the precipitate and thus, for small particles, with the disappearance of Guinier-Preston zones. Analysis of the results obtained in this research on the age-hardening of Cu-Be alloys furnishes additional evidence in support of this point of view.

Since the data on Guinier-Preston zone formation were obtained using a single crystal specimen it is desirable that the hardness values of the same specimen be used in studying the correlation of Guinier-

Preston zone formation and hardness. In Fig 11 it is seen that the hardness of the single crystal increased steadily on aging a t 200°C for 1,000 hr and correspondingly the Guinier-Preston streaks were growing in intensity during the same period, Fig 6. On aging a t 300°C the hardness maximum was reached in about 50 hr, Fig 12, and the Guinier-Preston streaks were observed to increase in prominence during this period. The initial softening at this tem- perature was accompanied by the begin- ning of the formation of X ray diffraction spots from the equilibrium precipitate after IOO hr aging. At 4oo°C the hardness maximum occurred in 4 hr, Fig 13; by this time the Guinier-Preston streaks were found to be well developed and there was evidence that the formation of equilibrium precipitate was about to begin as shown by the incipient condensation of the streaks into ill-defined spots.

No transition lattice was detected in the Cu-Be system. Coherency between pre- cipitate and matrix implies that the coherent precipitate exhibit either a lattice quite different from the equilibrium pre- cipitate, with dimensions adjusted to that of the matrix thus affording coherency (as in A1-Cu), or a lattice of the same type of the equilibrium precipitate merely strained in dimensions so as to afford coherency (as in A1-Ag). The coherent precipitate lattice in either case may be designated as a transition lattice but the proof of the occurrence of such a lattice demands that the precipitate particles grow to such a large size whilst remaining coherent that three-dimensional diffraction be possible. Presumably in the Cu-Be system the transition lattice breaks away from the matrix lattice (loses coherency) when the particles are too small to afford three-dimensional diffraction; other sys- tems show this behavior.20

The results obtained on grain boundary precipitation strongly suggest that precipi- tation in this region is similar to that

22 &tECHANISM OF PRECIPITATJON IN ALLOYS OF BERYLLIUM IN COPPER

occurring in the grain center, except that aging and overaging is accomplished more rapidly in the vicinity of most grain boundaries. The cause of this acceleration of the precipitation reaction may be attributed to the stresses that are present at the boundary between two crystals having different orientations.

I. The technique of quenching from the solution temperature directly to the aging temperature was applied to the study of the precipitation hardening of Cu-Be alloys. Little difference was observed in most properties compared with those obtained after water quenching and aging.

2. The change of electrical resistance during aging was observed in single crystals and in polycrystalline specimens. The single crystals reacted more slowly than did the polycrystals. At lower aging temperatures the anomaly in electrical re- sistance previously reported was confirmed.

3. In addition to the usual hardness tests carried out on single crystals and on poly- crystals, the Tukon hardness tester was used to measure the hardness of the grain- center and the grain-boundary areas of tri-crystals. I t was shown that the grain- boundary area hardens more quickly but also overages sooner.

4. Laue X ray patterns were used to demonstrate the existence of Guinier- Preston zones in Cu-Be alloys. An analysis of these patterns showed that these zones form on the (roo] planes of the matrix.

5. An improved Debye X ray technique involving vacuum operation and a mono- chromated beam failed to disclose a transi- tion lattice for this system. I t was concluded that no lattice other than that of the equi- librium precipitate grows large enough to afford three-dimensional diffraction.

6. Excellent evidence for the accelera- tion of hardening by the strains produced by water-quenching was obtained.

7. In studying tri-crystals evidence was

found to indicate that strain is the cause of grain boundary precipitation.

8. It i s concluded that the precipitation- hardening behavior of Cu-Be alloys can be explained by the theory that has been used to account for hardening in Al-Cu and Al-Ag alloys.

ACKNOWLEDGMENTS * Mr. Leon Fletcher of the Brush Beryl-

lium Co. very kindly supplied the copper- beryllium master alloy as well as chemical analyses of the iinal alloys.

Some of the results reported in this paper were obtained by the first author while he was employed in the Research Laboratory of the General Electric Co.

This research was done on a fellow- ship granted by the Metals Research Laboratory of the Carnegie Institute of Technology.

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A. G . GUY, C. S. BARRETT AND R. F. ME=-TP 2341 23

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