The effect of heat-treatment on the mechanical properties of some commercial beryllium

266

JOURNAL

THE

EFFECT

OF THE LESS-COMMON

OF HEAT-TREATMENT

PROPERTIES

OF SOME A. W. JONES

Research

and Development

(Original manuscript

Department,

ON THE

COMMERCIAL AND

MECHANICAL BERYLLIUM

R. T. WEINER Company,

United Power

received February

METALS

Erith,

z3rd, 1963; revised manuscript

Kent

(Great

Britain)

May zSth, 1963)

SUMMARY The effect of heat-treatment in the range 650°-950°C for times between halfan-hour and 500 h on the tensile properties under fast strain rates at 6oo’C of a batch of hot-pressed and warm-extruded virgin Brush QMV beryllium has been studied. The creep properties of this material were also investigated in the range 500”-6ooOC in the as-extruded condition and after heat-treatments at 950°C and 700°C. Metallographic and X-ray techniques were used to follow changes of structure and constitution brought about by heat-treatment and by strain. Heat-treatment at low temperatures was found to increase the ductility at fast strain rates but decrease the creep strength, whilst heat-treatment at high temperatures was found to decrease the ductility at fast strain rates and increase the creep strength. Heat-treatment had no effect on tensile strength or on creep ductility. The results are attributed to a precipitation phenomenon involving intermetallic phases, which have been identified by X-ray diffraction methods, and to the presence of oxide particles within the grain boundaries.

INTRODUCTION

Recent

worki-

has shown that by suitable

heat-treatment

it is possible to eliminate

the ductility minimum which occurs at about 600°C in tensile tests under fast strain rates on beryllium. However, it has also been found5 that in tensile tests performed at slower strain rates the elevated temperature ductility trough remains, even in heat-treated beryllium. The purpose of this investigation was therefore to investigate further

the effect of heat-treatment

it is strained

on the mechanical

properties

of beryllium

when

at slow rates. EXPERIMENTAL

Material The beryllium used was obtained from the Brush Beryllium Corporation in the form of rods. These rods had been fabricated from one batch of Q.M.V. grade, 200 mesh virgin Brush powder by vacuum pressing between1050” and IIOO”~, the material being held in this range for 17 h and then extruded at a temperature between 420’ and 450°C with an extrusion ratio of 6 : I The resultant rods were subjected to a postextrusion anneal of 15 min between 750~ and 775°C by the manufacturerbefore delivery. The results of a chemical analysis are given in Table I, together with the manufacturers specification. J. Less-Common

Metals, 6 (1964) 266-282

EFFECT OF HEAT-TREATMEST TABLE THE

MAKER’S

SPECIFICATION

AND

THE

OS Be

267

I

CHEMICAL

.IN.U_YSIS

OF

THE

MATERIAL

Muker’s speclfm&nt

(wt.

_

“6)

9s.7 1.35

Assay R,( )

oxygen

Carbon

Specimen

0.04

0.O.j

0.O.j

0.O.j

0.00

0.07

Silicon

0.031

0.03’

“.“JL

IroIl

0.16

0.13 0.13

0.13 0.13

Xluminium

0.047

o.oq

0.O.j

Magnesium

0.029

0.06

0.04

Density

I ,853 g/cm3

preparation

Hounsfield I

0.8 (1.3 H,O)

No.

0.13 cl.12

and heat-treatment

12

specimens

were used in all tensile

tests,

and specimens

with a

in. gauge length and 0.05 in.2 gauge area were used for all creep tests. The usual precautions

and cracking;

further,

were machined solution

0.010

containing

Heat-treatments to 950%

were taken during machining

to minimise

surface cold work

to ensure freedom from surface cold work the creep specimens in. oversize,

the excess metal being removed

5% CrOs, 83% HZ04

and

were carried out at 50°C intervals

Times from o to 500 h were employed

by etching

in a

H&O4 at 90°C.

12%

in the temperature

range 650”

in the range 650” to 85o”C, but at

900” and 950°C times were restricted to a maximum of 50 and 8. h, respectively. All heat-treatments were performed under IO p.s.i.g. of argon in controlled atmosphere furnaces

with the specimens

wrapped in tantalum

foil. At the end of the heat-treat-

ment the specimens were cooled to below 500°C in less than three minutes unless a slow cool was required in which case the temperature controller was programmed to reduce the furnace

temperature

at 5” Cjh.

Tensile and creep testing procedure The tensile testing machine used was substantially MORROW

tensile mercury

AND

MARTINI

tests were carried and at a strain

the tensile

curve

being

out under a dynamic rate of

1.8qg

the same as that used by autographically

vacuum

of better

per min. As-extruded

recorded. than

material

10-3

BROWN,

All the mm of

was tested

at

400’, 500”, 600~ and 7oo”C, but tests on heat-treated specimens were restricted to 600°C. The creep frames used were of modified B.N.F. design and were fitted with gas vessels to enable all the testing to be carried out in argon. Creep tests were performed at 5o0°, 550” and 600°C. Metallography One half of each metallographic blank was heat-treated concurrently with the relevant tensile or creep specimen whilst the other half was retained as a standard.

268

A. W. JONES,

R. T. WEINER

The fractured tensile specimens which were examined metallographically were mounted without sectioning and were then polished back to give a longitudinal section. This technique was also employed on portions taken from the gauge length of fractured creep specimens. The remaining portions, including the heads, were sectioned along the axis of the specimen. All specimens were mounted and polished in the normal manner and examined under both bright field illumination and polarised light. Grain sizes were measured in both the longitudinal and transverse directions using polarised light. Electron microscopy Electron microscopy was used to examine residues obtained when specimens were dissolved in bromine-methanol solutions. The residues were centrifuged onto glass slides, which had been covered on one side with an evaporated carbon film. The films were then floated off and collected on standard grids. Where selected area diffraction was considered desirable the mounted specimens were shadowed with gold to provide an accurate calibration. X-ray diffraction A number of specimens with different thermal and strain histories were examined by means of X-ray diffraction techniques. Conventional 11.46 cm and rg cm powder X-ray cameras were used, together with nickel-filtered copper radiation or ironfiltered cobalt radiation. In general, X-ray patterns were produced by a glancingangle technique using solid specimens. Residues from the bromine methanol extractions were also examined by X-ray diffraction methods. RESULTS

Tensile tests The variation of the tensile properties of as-extruded material with testing temperatures is shown in Figs. Ia, Ib and IC. The limit of proportionality was taken as the stress at which the load extension curve first exhibited deviation from linearity. In spite of considerable scatter it may be seen that both the strength and ductility of the material decreased with increasing testing temperature. The tensile properties after various heat-treatments are shown in Fig. za-d. Lines have been added to these histograms indicating the range of results obtained from each heat-treatment temperature. It can be seen that the strength parameters have skew distributions with pronounced maxima, and do not appear to depend on the heat-treatment temperature. The ductility on the other hand appears to increase with decreasing heat-treatment temperature. Statistical analysis of the results confirmed these conclusions, the effect of heattreatment temperature on the ductility being significant at the 0.1 yO level. Although the effect of heat-treatment time on ductility was significant at the 5% level, the correlation coefficient between ductility and time, with effect of temperature eliminated, was +O.ZI, a value which is statistically, not significantly greater than zero. The apparent effect of time was probably due to the fact that it was not possible to balance the experiment statistically and that there was therefore an association between high temperature and short times, and between low temperature and long times. J. Less-CommonMetals,

6 (1964) 266-282

ON Be

EFFECT OF HEAT-TREATMENT

t.0

d q

rl

0

L.Of P. Fracture

5tress

30

Fig. IC. Elongation UPYSUS testing temperature for asextruded material.

Fig. la. U.T.S. and L. ofP. U~YSUStesting temperature for as-extruded material.

1

400

LOO

500

7

1

Temp.(‘YI

Fig. lb. Fracture stress and I,. of P. uwsus testing temperature for as-extruded material.

Specimens

cooled slowly from the heat-treatment

to specimens cooled rapidly. No multiple yielding of the type described

temperature

gave similar results

by other workersl,G

was encountered

in any of the tests. All fractures

were intercrystalline,

and only the specimens

tested at 400’ and SOOT

showed necking.

r

16 L. of

Fig. *a. Histogram

P. psi.

18

20

22

24

26

28

30

x 10J

of L, of P. Results obtained in tensile tests at 600°C on heat-treated j.

Less-Comnzon

Metals, 6

(1964)

material. 266-282

A. W. JONES, R. T. WEINER

270 0-

-650T-

c-

_6OO~C

I ,~QOOTB

7‘

L15extrude-

I-7OOT-l

750°C 650-C-

-

-

I-950’C-l

6-

ul 5z 2

4-

%

1

g!._ I__ 0

2

4

6

0

10

12

14

U.TS

Fig. zb. Histogram

16

psf.xlOJ

18

20

22

24

26

28

30

of U.T.S. Results obtained in tensile tests at 6oo’C on heat-treated

-As -65O”C-

8.

, ~sOO”C~

7-

material.

,

extruded

c7OO~C+

c75ooc~800°C~ -5ov-9

6 I

0

2

4

6

8

10

12 Fracture

Fig. PC. Histogram

14

16 stress

of fracture stress results obtained materials.

18

20

22

24

26

28

30

p.s.i.x IO3

in tensile tests at 600°C on heat-treated

5 VI f Q4 B .3 z 2

1 I 0

2

4

6

0

10

12 01. elongation

Fig. 2d. Histogram

1 18

20

22

24

26

28

30

of ductility results obtained in tensile tests at 6oo’C on heat-treated J. Less-Common

Metals,

6

(1~64)

materials. 266-282

EFFECT OF HEAT-TREATMENT TABLE CREEP

AND

RUPTURE

DATA

i”Ci 3,500 3,000 2.500 2,250

None

200

h at 700%

24 h at 95o’C

AND

9.8 1 1 1.6 0.8

600

0.54

I,jOO

0.2

4,000 31500 3,000 3.000

600

weep

x 10--s (ita.jin /iI)

Yale

2,000

4,000 2,500 1,500 1,000

6,000

550

5,000 4.000

I.7 0.9 I.4 2.1

12s

1.4

1351

I.1

3.0 I.3

123 374 1402

3.7 3.4 4.7

190 25 7.5

80

2.6

295 422

0.7 3.4

I.45 0.55

774 I229

2.4 3.1

0.21

I749

2.1

96

2.4 3.9

10,000 None

zoo

9,000 8,000 h at 700°C

IO,000 9.000 8,000

._.-

99 2.7 2.4 324 2.4 479 1093 1.3 Test discontinued Test discontinued

‘35 ‘67 661 1278

4.9

14.0

200 h 700°C.

lhJ ---___-

6.2

2,000

550

tame

3.1 2.7 I.9 1.1

0.86

None

Rupture

5 72 ‘87 578

2,000

6,500

BERYLLIUM

456 49-5 4.3 2.4

600

5*500 5,000

271

HEAT-TREATED

Minimum

Temp.

Be

II

ON AS-EXTRUDED

Ttist conditions -

ON

500

20 j.1

500

2.9

410 422

2.1

Creep tests From the tensile results it became clear that there were twostatesof heat-treatment, one characterised by a low ductility and obtained by heat-treatment at a temperature above 850°C, the other characterised by a high ductility and obtained by heattreatment at temperature below 800°C. In order to investigate the creep properties of both these states tests were carried out on material heat-treated for 24 h at 950°C and on material heat-treated for 200 h at 7oo’C. The results of these tests together with those on as-extruded material are shown in Table II. Both in terms of minimum creep rate and rupture time the material heat-treated at 950°C showed the greatest creep strength and the material heat-treated at 700°C showed the least, with the strength of the as-extruded material in an intermediate position. The creep ductility was low in all cases, however, and did not appear to depend either on the heat-treatment, testing temperature, or testing time. All fractures were intercrystalline in nature and the specimens showed no necking. The /.

Less-Common

Metals,

6 (1964)

266-282

A. W. JONES, R. T. WEINER

272

amounts of primary creep strain that occurred were least in specimens heat-treated at 950°C and most in specimens heat-treated at 7oo”C, with the as-extruded material again in an intermediate position. Metallography

The as-extruded material had an average grain size of between 0.020 and 0.025 mm and this was not increased by heat-treatments of 50 h at goo”C, 24 h at 95o”C, or 2 h at IOOO’C. Short heat-treatments at temperatures above IOOO’C, or longer heattreatments at ~ooo”C, did however cause an increase in grain size. The grains were found to be slightly elongated in the direction of extrusion. Table III shows the variation of grain size for tensile specimens in the as-extruded condition and in two TABLE THE

GRAIN

SIZES

OF SELECTED

Heat-treatment Temp. (“Cl

700

AFTER

TENSILE

Grain count (gvainslmm)

Time

Longitudinal -

TESTING

AT

600°C

Average grain

Transverse

diameter (mm)

(h)

500

As-extruded

950

SPECIMENS

III

8

Grip 46.75 Gauge 37.2

49.0

0.0205 0.0239

Grip 45.0 Gauge 41.75

46.5 48.9

0.0218 0.0221

Grip 44.25 Gauge 42.02

46.75 49.75

0.0218

51.0

0.0220

Fig. 3a. Grain structure of tensile specimen heated-treated at 700°C for 500 h and fractured 600% Polarised light, etched in I % HF in glycol. (x 150) J. Less-Common

at

Metals, 6 (1964) 266-282

EFFECT

OF HE.\T-TRE.iThlENT

ON &

273

Taking the grain size in the grip portion of the spy:cimen con ditions of heat-treatment. condition it may be seen that more grain elon gation asc zharacteristic of the unstrained occ urred in the specimen treated at 700°C than in specimens in the as-extrude d conditi on or heat-treated at 950°C.

Fig, 3b. Grain structure

Fig. 3~. Grain structure

of as-extruded in

tensile specimen fractured

I “h HI; in gly~A. (x 150)

at 6oo”C. Polarised light, etched

of tensile specimen heat-treated at 950°C for 8 h and fractured Polarised light, etched in I “: I-IF in glycol. ( x 150) J. Le.v-Comnwn

M&Es,

at Goo”C.

6 (1964) 266-282

274

A. W. JONES, R. T. WEINER

Fig. 4a. Bright field micrograph showing the fracture area of a creep specimen heat-treated at 700’ ‘C for 200 h and tested at tio”C (fracture time 72 h). Specimen etched in 1% HF in glyc :ol , (x50)

Fig. 4b. Bright field micrograph showing the fracture area of creep specimen heat-treated at 950°C for 24 hand tested at 6oo’C (fracture time 135 h). Specimenetchedin I yQ HF inglycol. ( x 50) J. Less-Common

Metals, 6 (1964) 266-282

EFFECT OF HE.\T-TREATMENT OX k3

275

Fig. qc. Bright field micrograph showing the fracture area of an a>-vutruded creep specimen tested at 600°C (fracture time gg h). Specimen ctchrd in I “(, tIF in glycol. ( x 50)

Intercrystalline cracks were observed in the gauge portions of all the specimens that were examined, but in the case of the specimens heat-treated at 700°C intercrystalline

cracking

of sections

taken from the gauge lengths

was restricted

to the area near the fracture. of fractured

Typical

tensile specimens

micrographs are shown in

Figs. 3a, b and c. The white spots which may be seen in these micrographs precipitated particles. The creep specimens fractured in about up of intergranular

shown in Figs.

qa, b and c were all tested

are due to

at 600°C and all

IOO h. It was found

cavities.

that failure had taken place by the linking There was a marked difference in the amount of cavita-

tion that had taken place, the specimen heat-treated at 700°C showing the most extensive cavitation (Fig. 4a), whilst there was considerably less in the specimen heattreated

at 950°C (Fig. 4b) and least in the specimen

creep tested in the as-extruded

condition. No grain growth was found in any of these specimens. It can be seen from Fig. qa that fracture of this specimen was associated with a defect. It seems possible that such defects were in part responsible for the considerable experimental

scatter

found in the results of the tensile

and creep tests.

X-ray and electron microsco~!Gc observations In all the specimens examined, irrespective of heat-treatment, small amounts of beryllium oxide, beryllium carbide, free silicon, and an intermetallic phase designated Bei3X were found. X-ray and electron microscopic investigation of extracted residues showed the oxide particle

size to be independent

the range 0.1-1 micron. The intermetallic compound

Bei3X

of heat-treatment

and to be within

was in the form of large, scattered

particles

of

J. Less-Common Metals, 6 (1964) 266-282

950

g o\ P

9oo

950

5

z !, :

24

950

& ” 0

2

24

8

950

Vacuum

Argon

M.

19

700

24

24 24

7oo 750

M. 17 M. 18

24

7oo

Vacuum

24 24 24 24 24 24 24 -

430 500 600 7oo 750 800 900 -

-

13 74 14

12

7 8 9 IO II

24

SPECIMENS

-

c.

I

-

24

Tested for 1210 h at 6oo’C

-

700 T. 3

-

24

24

700

750

-

-

hf. 15*

2

24

7oo

-

(h)

Time

Strained specimens Temp. W)

M. 15

T.

T. I T. I*

M. 3 M. 3*

Specimen no.

Second heat treatment

A(S) and B(V.W.)

B B

None

None None B(W) B(V.S.) B(S) B(S) None None B

B

None A(S)

B(S) B(S)

Phases

B(S)

24

IV HEAT-TREATED

‘4

M. 16

M. M. M. M. M. M. M. M. M.

TABLE VARIOUS

-

24

20

(h)

Time

Argon

Vacuum

Vacuum

24

1000

8

750

M. 6

rZrgon

24

1000

3 s

2

R ‘:

b

4

750

M. s M. 5

Argon

750

2

1000

4

Vacuum

24

1100

M.

Argon

2

7oo 750

Temp. (“C)

1100

M. I M. 2

Specimen no.

IN

Unstrained specimens

Vacuum

Atmosphere

2

Time (h)

DETECTED

1100

Temp. I”C)

First heat-treatment

PHASES

Head A(M) Gauge A(V.S.)

A

None A(S) and B(W)

A(S) and B(S)

None B(S)

None B(S)

Phases

160

soo

5oo

5oo

5oo

850

850

800

7oo

650

20

-

xl. 22

M. 23

argon

-

Argon -

.4(M)

A(S)

None l3

A(W)

None T. 4

c. 2

T. 8

T. 7

T. 6

24

1000

Argon

7oo

600

-

T. j

M. 21 M. 21

nf.

M. 20

Argon

Vacuum

Argon

-

-

-

-

-

C =

creep specimen;

Tested for 1093 h at 6oo’C

-1 = Re#e,Al); B = BQIZ (where Z probably is Fe); M = metallographic blank; T = tensile specimen; S = strong; M = medium; IV = weak. * Specimen strained at room temperature between first and second heat-treatments.

-

50

900

V.S.

=

xrery strong;

Head A(M) Gauge .4(V.S.)

.\(S) and 13 (Trace)

A(S)

A

x

A

F

2

3

2 m

$

2

5

8

-,

2

278

A. W. JONES, R. T. WEINER

up to 30 microns diameter. It was cubic and had a lattice spacing of about 10.4 A, and was therefore isostructural with BelsMg 7. As magnesium is used during the extraction of the metal, it seems likely that X is magnesium or a mixture of elements in which magnesium is prominent. Two further phases were detected, but the occurrence of these depended on heattreatment as shown in Table IV. The phase A was f.c.c. and was of the type Be5Y, with a lattice spacing of about 6.1 A. This was identified as Bes(Fe,Al)s. The phase B was hexagonal and was of the type Be&?. It was barely distinguishable from the BelIFe (or BelzFe) phase reported by TEITEL AND CoHENg. It can be seen from Table IV that heat-treatments above about 850°C took both phases into solution. Solution treatments for 2 h at IIOO’C or for longer times at lower temperatures, particularly in vacua, favoured precipitation of the phase B during the subsequent ageing heattreatment, whilst the application of strain before or during the latter heat-treatment favoured the precipitation of phase A. Material in the as-received condition was found to contain Bes(Fe,Al). This therefore suggests that precipitation occurred during the post extrusion anneal. As delivered, the material was in the aged or partially overaged condition. DISCUSSION Effect of heat-treatment

on tensile properties

It has been fairly well established 2- 499 that the elevated temperature ductility trough found in tensile tests at fast strain rates on beryllium is due to strain-induced precipitation of the phase A referred to above. The ageing treatment is thought to improve the ductility by precipitating this compound in the form of coarse particles which have little effect on the movement of dislocations. In the present investigation the X-ray diffraction studies and the improvement of ductility brought about by ageing below 850°C are in agreement with the proposed strain-induced precipitation mechanism. Further evidence in support of this mechanism is provided by the observation that the ductility of the material increased as the ageing temperature was reduced. However, several observations were not in agreement with this mechanism and these are listed below: (I) The lack of response of the U.T.S. and fracture stress to heat-treatment would not be expected. (2) Whilst some improvement of ductility was obtained by heat-treatment, contrary to the observations of other workers this was not accompanied by a change in fracture mode, all the fractures being intercrystalline. (3) If the strain-induced precipitation were the only embrittling mechanism then the ductility of the aged material at 600°C ( N 25% elongation) should not be less than the ductility of the as-extruded material at 400°C (N 37% elongation). (4) If strain-induced precipitation alone was taking place in the solution-treated material then the rate of work hardening would be expected to be greater for this condition of heat-treatment than for the aged condition. No such difference has been observed, the two load-extension curves shown in Fig. 5 being typical of the two states of heat-treatment. It is interesting to note that a considerable reduction of load occurred prior to fracture in spite of the fact that no necking was observed. It is therefore clear that some factor in addition to strain-induced ageing influenced the mechanical properties of the metal. J. Less-Common

Metals,

6 (1964) 266-282

EFFECT

Whilst

no direct evidence

was obtained, occurred

279

as to the nature of the deformation

it is probable

simultaneously

ON Be

OF HEAT-TREATMENT

that both grain boundary

during straining.

process taking place

sliding and grain deformation

If it is assumed

that the observed

inter-

crystalline cracking was initiated by the interaction of grain boundary particles and sliding grain boundaries then the independence of the tensile strength parameters at 600°C with respect to heat-treatment fact that grain boundary oxide particles ment.

may be explained as a consequence of the were present in all the states of heat-treat-

I

0

I

2

3

4

5

6

7

(I

9

lo

I,

12 13

84

Is

,e

16 0

Ip

20

II

22

23

a

25

26

7 28

Extenston(‘3~) Fig. 5. Tensilecurves for two statesof heat-treatment.

The fact that the ductility

of the aged material

at 600°C was less than that of as-

extruded material at 400°C may be explained if, as in aluminiumro, the relative amount of grain boundary sliding increases with increasing testing temperature. Such a change in the mode of deformation

would result in an increase

between sliding grain boundaries and grain boundary perature and a decrease in ductility. Effect

of heat-treatment

particles

in the interaction

with increasing

tem-

on creep behaviour

Whilst heat-treatment was found to affect the ductility but not the strength tests carried out at fast strain rates, it can be seen from Table II that in tests slow strain rates the creep strength

was affected

but the rupture

ductility

in at

was un-

altered by heat-treatment. COTTRELL~~ has suggested that cavities of the type observed in Figs. 4a, b and c can be nucleated by the interaction of sliding grain boundaries with non-wetted particles within the grain boundaries themselves. He has also shown that, once nucleated, the

cavities can grow under very small applied diffusion of vacancies. The X-ray diffraction of particles of BeS(Fe, Al) were present in the tion-treated material. It is therefore probable

stresses by a process of grain boundary studies have shown that a large number aged material but only a few in thesoluthat these particles, togetherwithoxide

280

A. W. JONES, R. T. WEINER

particles which are present in the grain boundaries of the material in all states of heat-treatment, were responsible for the cavities since relatively few were observed in the solution-treated condition but many more were present in the aged condition. The existence of different cavitation densities in creep specimens subjected to different heat-treatments prior to testing cannot, of itself, explain the difference observed in the amounts of primary creep and in the secondary creep rates since the nucleation and growth of cavities is thought to control only the onset and extent of the third stage of creep. Nearly all the proposed mechanisms for grain boundary sliding12 involve the movement of dislocations. It is therefore suggested that the strain-induced precipitation which was found to take place in solution-treated and as-extruded creep specimens (see Table IV), together with solid solution strengthening in the early parts of the tests, made the movement of dislocations more difficult and therefore reduced the rate of grain boundary sliding. In the solution-treated material the rate of grain boundary sliding would therefore be slow and the number of nuclei for cavitation would be small, and this would lead to long rupture times. On the other hand in the aged material the rate of grain boundary sliding would be fast and there would be a larger number of nuclei for cavitation. In this case rapid formation of cavities and early fracture would be expected. The mechanism of failure by the linking up of cavities clearly limits the ductility of beryllium under creep conditions. Due to the presence of oxide in thegrain boundary in all conditions of heat-treatment a situation favourable for the formation of voids exists in all cases, and no amelioration of creep ductility can be brought about in powder fabricated beryllium by heat-treatment. OLDS et aZ.5 came to a similar conclusion on the basis of tensile tests carried out at constant slow strain rates. X-ray

diffraction

studies

Whilst several other workers 2,4,1s have found that the phase Bes(Fe,Al) was the precipitating phase in commercial beryllium, only POINTU et al.4 have reported finding a phase of the type BeizM, where M is a transition metal; unfortunately they do not give the conditions under which it was found. It seems possible that the Be& phase was not found by the other workers because the heat-treatments which they applied were all at lower temperatures than those applied to some of the specimens investigated here. MOORE et al.2 have pointed out that from a consideration of the beryllium-iron equilibrium diagram a precipitate of the type BeliFe rather than Be5Fe would be expected. Electron microprobe investigations 2,3,13 have shown that particles rich in aluminium and silicon are present in commercial beryllium, and MEREDITH et al.3 found that during ageing iron precipitates onto these particles. The present investigation indicates that, although the precipitating phase may be BeliFe, association with aluminium causes a transformation to the Bes(Fe,Al) phase. Heat-treatment at about 850°C apparently therefore takes only the precipitated iron-containing particles into solution, leaving the aluminium and silicon-rich phase unchanged. The solubilities of aluminium and silicon in beryllium are 10~14, and since the metal investigated here contained less than 500 p.p.m. of each it is possible that heattreatments at higher temperatures for long times took all three elements into solid solution. If aluminium is taken into solution, then on ageing the iron and aluminium J. Less-Common

Metals,

6 (1964) 266-282

EFFECT OF HEAT-TREATMENTON Be might precipitate Bes(Fe,Al)

separately

and the aluminium

would not be available,

necessary

leading to the precipitation

281 for the stabilisation of a BelIFe-type

of

phase.

MOORE et ~1.2have shown that the elevated-temperature ductility of a berylliumaluminium alloy can be improved by ageing, and they conclude from this that a strain-induced ageing process takes place in the alloy. This may explain why strain favoured

the precipitation

possible that aluminium

of BeS(Fe,Al)

in the present investigation,

for it seems

and iron could then become associated by being precipitated

on dislocations. The possibility of aluminium loss during the solution heat-treatments was investigated but no positive evidence of such a change in composition was found.

CONCLlJSIONS (I) A phase Bes(Fe,Al) powder fabricated

was found by X-ray diffraction

beryllium

by heat-treatments

in the as-extruded

at temperatures

condition.

studies in some commercial This phase was dissolved

of 850°C and above. In certain circumstances

it

was reprecipitated on subsequent ageing below 850°C. (2) Heat-treatment at or above 85o’C produced a small decrease in the tensile ductility

of this material under fast strain rates at 600°C. This is attributed

to the

strain-induced precipitation of the dissolved Bej(Fe,Al) during testing. (3) Heat-treatment below 850°C produced a small improvement in the tensile ductility under fast strain rates at 600°C. This improvement tion of strain ageing during the tensile tests. (4) Heat-treatment

is attributed to the reduc-

did not appear to have any effect on the limit of proportionality,

on the ultimate tensile stress, or on the fracture stress. These parameters are thought to depend on the interaction of sliding grain boundaries with beryllium oxide, the amount and distribution

of which was independent

of heat-treatment.

(5) The creep ductility was also independent of heat-treatment. this was again due to the interaction of sliding grain boundaries particles which led to cavitational

It is concluded that and grain boundary

failures after limited extensions.

The variation

cavity densities appeared to depend upon the number of Fes(Fe,Al) in each state of heat-treatment.

particles

in

present

(6) When compared with the creep strength of material intheas-extrudedcondition, the creep strength was improved by heat-treatment at 950°C and reduced by heattreatment at 700°C. These changes are attributed to the effect precipitation hardening on the rate of grain boundary sliding.

of strain-induced

(7) In the absence of strain Bes(Fe,Al) was only precipitated when the solution treatment was of short duration at temperatures below IIOO’C. After heat-treatment at IIOO’C for two hours, or for longer times at lower temperatures, vucuo, a phase Bell2 (nearly identical to BelIFe) was precipitated effect of this phase on mechanical properties was not investigated.

particularly

in

on ageing. The

ACKNOU’LEDGEMENTS The authors would like to thank Mr. A. TRINER for the statistical analysis of the tensile results and Miss S.R. BILLINGTONfor the metallographic examinations. Special thanks are due to Mr. H. P. ROOKSBY of the Chemistry and Technical Services J. Less-Common Metals, 6 (1964) 266-282

282

A. W. JONES, R. T. WEINER

Division, Central Research Laboratories, The General Electric Company Ltd., Hirst Research Centre, Wembley, for X-ray diffraction measurements, interpretation and many valuable discussions. Financial support for part of this investigation was obtained under extra-mural contract from the U.K.A.E.A., and this paper is published by permission of the Managing Director of the Reactor Group of the Authority. REFERENCES 1 A. R. BROWN, F. MORROW AND A. J. MARTIN, J. Less-Common Metals, 3 (1961) 162. 2 A. MOORE, F. MORROW, V. D. SCOTT AND D. A. CHEER, Conf. on the Metallurgy of Beryllium, London, 1961. 3 J, D. MEREDITH AND J. SAWKILL, Conf. on the Metallurgy of Beryllium, London, 1961. 4 P. POINTU, L. ESPAGNO, L. Azou AND P. BASTIEN, Compt. Rend., 250 (1960) 2365. 5 G. C. E. OLDS, T. RAINE, J. A. ROBINSON AND A. G. TODD, Conf. on the Metallurgy of Beryllium, London, 1961. 6 P. BASTIEN AND P. POINTU, J. Nucl. Mater., 4 (1961) 113. 7 T. W. BAKER AND J. W. WILLIAMS, Acta Cryst., 8 (1955) 519. * H. P. ROOKSBY, ,I. Nucl. Mater., 7 (1962) 205. 9 R. TEITEL AND M. COHEN, Trans. AIME, 185 (1949) 285. 10 D. MCLEAN AND R. C. GIFKINS, J. Inst. Metals, 8g (1960) 29. 11 A. H. COTTRELL, Syffap. on Structural Processes in Creep, Iron and Steel institute, London, 1961, p. I. 12 R. C. GIFKINS, J. Australian

Inst. Metals, I (1956) 134; and Fracture, Wiley, New York, 1959. Ch. 27. 13 A. K. WOLFF, S. H. GELLES AND L. R. ARONIN, Conf. on the Metallurgy of Beryllium, London, 1961. 14 G. E. DARWIN AND J. H. BUDDERY, Beryllium, Butterworths, London, 1960, Ch. g. J. Less-Common Metals, 6 (1964) 266-282