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The effect of thermal treatment on the metallographic structure of beryllium and dilute beryllium alloys
JOURNAL
THE
OF NUCLEAR
EFFECT
~~
lo,2
OF TEENY
(1963)
TREA~ENT
BER~LI~
~OR~-EO~~D
113426,
ON THE ~ETALLOGRAPHIC
AND DILUTE
BERYLLIUM
C. V. T. RANZETTA Metallurgy
.lXv&iorc, Atomic
Wmpons
PUBLIS~G
CO.,
AMSTE~~
STRUCTURE
OF
ALLOYS
and V. D. SCOTT
Research Establiahmmt,
Akkrmmtm,
Be&a.,
UK
Received 13 March 1963
Electron microscopy and electron-probemicroanalysis techniqueshave been used to study t,hemiorostructure of beryllium and some dilute beryllium alloys. The effect of thermal treatment on the dissemination of impuritiesand alloying elements has been investigated with particular referenceto the mechanical properties of the metal at 600” C. In beryllium-iron alloys annealed to improve hightemperature ductility the structure was typified by the presence of large iron-containing grain-boundary precipitates, together with a more general segregation of iron to grain boundariesand other favourable sites. Some indication of solid-solution hardening in this system was apparentbut no evidenceof a precipitationageing process was found. It is therefore suggested that the microstruct.urein the grain-boundary regions of the metal may be important in affecting the ductility and coherency of metal grains at high temperatures. Annealed beryllium-ohromium alloys had good high-temperatureductility which in this case may be associated also with the observed purification of the matrix.
La microscopic Blectronique et la technique de micro-analyse par sonde Blectroniqueont Bte utilisdes pour Studier la microstructure du b&yllium et de quelques alliages peu oharg& en Elements d’addition de beryllium. L’influenoe du traitement thermique sur la repartition des impuretes et des Elements d’alliage a Bte t%udit%een liaison Btroite avec les propri&tb mecaniques du m&al it 600” C. Dans les alliages b&yllium-fer, recuits dans le but d’ameliorer la duotilite it haute temperature, la structure Btait caract&isee par la presenoe de gros precipites riches en fer aux joints de grains ainsi que par une segregation plus g&&ale du fer, soit aux joints de grains, soit au hasard. On a mis en 6vidence
un certain durcissement structural dans cet alliage, mais on n’a pas trouve de preuve caraoteristiquedu processus de vieillissement aocompagne de precipitations. C’est pourquoi on a suggere que la, ductilitb du metal et la cohesion entre grains it temperature &e&e devraient d$pendre $troitement de la mierostructure existant au voisinage des contours de grams. Les alliages Be-Cr recuits possedent une bonne ductilite a haute temp&ature, fait qui peut Gtre Bgalementrelit5&la purificationobserveede la matrice.
Es wurden elektrone~~oskopische und elektronenm~o~al~~che Verfahren zur ~ntersueh~g des Gefiigesvon Berylliumund einigenl&lichen BerylliumLegierungen benutzt. Der Einfluss der Warmebehandlung auf die Verteilung von Verunreinigungen und Legierungselementen wurde untersucht, wobei besonders auf die mechanisohen Eigenschaften des Metalls bei 600” C geachtet wurde. In Beryllium-Eisen-Legierungen, welche zwecks Verbessernngihrer Hoch~mperaturdu~ilit~t warmebehandelt wurden, war die Struktur typisiert durch die Gegenwartvon grosseneisenhaltigenKorngrenzenniederschlagenund einer mehr allgemeinen Abscheidung von Eisen an den Korngrenzen und anderen Stellen. Einige Anzeichen von Mischkristallh&tung in diesem System wurden festgestellt. Es konnte jedoch keine Klarheit iiber einen Niederschlagalterungsprozess gefunden werden. Es wurde daher ~geno~en, dass das Gefiige in den Korngrenzemegionen von wesentlichem Einfiuss auf die Dehnbarkeit und Koharenz der Metallkorngrenzen bei hohen Temperaturen ist. Die warmebehandelten Beryllium-Chrom-Legierungen hatten gute Hochtemperaturduktilitiit, was in diesem Fall such zusammenbangen mag mit der beobachteten Reinheit der Matrix.
113
114 1.
G. V. T. RANZETTA AND V. D. SCOTT Introduction
A considerable amount of evidence has been accumulated during the past few years showing that the mechanical properties of beryllium in the temperature range 400-600” C depend strongly on the composition and thermal history of the metal. It may be said generally that rapid cooling of commercial-purity beryllium or dilute beryllium alloys from high temperatures, such as those used for extrusion (- 1050’ C), leads to a fall in ductility with increase in temperature above 400’ C, to a ~~rnurn at about 600” C. This drop in ductility can be avoided by a slow cooling or annealing treatment which is usually more effective when applied to dilute beryllium alloys rather than to commercial-purity beryllium l). In the present paper, results obtained by the application of electron-probe microanalysis and electron microscopy to the study of such effects in beryllium and beryllium alloys are presented. Since tensile test data were available for most of the specimens examined it has been possible to correlate microstructural information with mechanical behaviour and to discuss the results in relation to previous theories proposed to explain this ductility phenomenon.
2. 2.1.
Experimental Details ORHXN OF MATERIAL
The bery~i~ used in this study was produced from Pechiney commercial-purity electrolytic flake which was melted, together with any desired additions, cast, and subsequently hotextruded to a size suitable for a mechanical evaluation of its properties at 600* C. In addition a range of binary and ternary alloys was made by arc-melting together small quantities of powdered flake beryllium and the required alloying element. This group of specimens was intended solely for microanalytical studies and was in a form u~uitable for mechanical testing.
ELECTRON-PROBE MICROANALYSIS
2.2.
Microanalysis was carried out using a Cambridge scanning electron probe inst~ment, quantitative measurements being made in most cases by reference to pure metal standards. 2.3.
ELECTRONBIICROSCOPY
Replica methods have been used to study the surface of both metallographioally-prepared and fractured specimens of a beryllium0.15 wt y0 iron alloy. In addition, thin foils of this metal have been prepared by chemical and electroche~cal dissolution for direct examination in the electron microscope. 3.
Results
To avoid the unnecessa~ repetition 03 thermal treatment details, specimens heat treated to give better ductility at 600° C will be referred to as “annealed” and those quenched from high temperatures, giving a less ductile product, as “homogenised”. 3.1.
ELECTRON-PROBE MICROANALYSIS
In general the low level of impurities in commercial-purity beryllium and beryllium-iron alloys has made the study of their dissemination ~fficult and dilute alloys consisting of bery~um and certain of the commonly-occnrring impurity elements have therefore been examined. With regard to quantitative microanalysis the volume of material irradiated by the electron probe will be about 30 prna depending upon analysis conditions. As a consequence the probe values quoted for inclusions smaller than about 4 pm in diameter may tend to be low; this size effect could be considerable for particles with a diameter of less than 2 ,um. X1.1.
Beryrllium- Iron Alloys
Alloys containing iron in the range 0.4 wt y. to a nominal 0.03 wt %, as found in commercial grade beryllium, have been investigated in both the “annealed” (ductile) and “homogeni~’ (brittle) condition.
~~TALLO~RA~H~C
STRUCTURE
OF
BERYLLIUM
s.
Fig. 1.
b.
Electron.
d.
X-my,
Ax.
1400 x
1400 x
Optical.
AND
DILUTE
BERYLLIUM
ALLOYS
400 X
0.
X-my,
Fe.
1400 X
8.
X-my,
Si.
1400 X
Beryllium - 0.18 wt % iron alloy, annealed at 800’ G for 120 houra and slowly cooled, &owing segregation ati gmin boundaries.
1x6
116
G.
a.
Fig. 2.
Electron.
Beryllium
V.
T.
RANZETTA
‘1000 x -
AnneaM
V.
D.
SCOTT
b.
X-ray,
Al.
0.15 wt y0 iron alloy, homogenised at 1000” C for quenched, showing alumiuium at grain boundaries.
Some preliminrtry results obtdned on alloys containing greater amounts of iron (1 wt */*and 5 wt %) sre also presented, 3.1.1.1.
AND
specimens (Ductile at
600” C)
A general enrichment of the grain boundaries in iron was found in these specimens. This was most pronounced in the higher iron alloys, 0.8 wt y. Fe being indicated at the grain boundary of a 0.4 wt y. alloy, but was still just deteotable in a 0.08 wt y. specimen. The grain boundary segregation of iron is clearly shown in figs. 1 a-c, which were obtained from a 0.18 wt y. iron alloy that had been annealed at 600’ C for 120 hours and slowly cooled. The degree of iron enrichment increased with time during heat-treatment, indicating that iron is slowly diffusing to the grain boundaries during annealing. For example, in the case of a 0.16 wt o/0 iron alloy heat-treated at 800” C, no iron was found after 2 hours, slight enrichment w&s visible after 10 hours, and after 100 hours the iron WM readily detectable. P&icles 1 to 3 ,.umin size and located mainly at grain bound~ies were fairly common in annealed beryllium specimens, ~ero~n~lysis showing the presence of iron in concentrations of up to 7 wt Oh, usually together with a few per cent aluminium and silicon (figs. 1 d-e).
1000 x 2 hours
and oil
A slight enrichment of the grain boundaries in aluminium may be observed which appears to be associated with the ~stribution of iron, but not every grain boundary was enriched to the same extent. The iron con~nt of the matrix was not found to be significantly affected by the heat treatment; a very small change ( < 0.02 wt %) would be within the experimental scatter of most of the measurements and remain undetected. 3.1.1.2. Homogenised specimens (Low ductility at 600’ C) Specimens of as-extruded and annealed beryllium which h&d been subsequently homoge~sed at high tem~r&t~es (90~1060’ C) and quenched &ve been examined. No grain-boundary enrichment was detected, the iron being distributed uniformly throughout the specimen irrespeotive of its initial dissemination; thus the iron segregation at the grain boundaries associated with increased ductility a;t 600’ C is reversible. Figs. 2a, and 2b were taken from it, grain boundary showing a, rela;tively high concentration of alu~~um-rich inclusions in a commerci&l-purity beryllium specimen that was oil-quenched after heat treatment at 1000” C. Up to 60 wt y0 aluminium was measured in these inclusions, no other elements being detected.
METALLOGRAPHIC
STRUCTURE
OF
BERYLLIUM
3.1.1.3. Specimens in an intermediate condition During extrusion and cooling, beryllium specimens are subjected to a form of heat treatment that ~a~ot be closely defined and it might therefore be expected that as-extruded specimens oould be in a condition in~~e~ate between the homogenised and annealed state. However, microanalysis results from such alloys were similar to those obtained for homogenised metal. No segregation of iron to the grain boundaries was observed but it was possible to locate inclusions slightly richer in iron than the matrix ; for example 0.3 wt $.&Fe was measured in an inclusion contained in an 0.18 wt y0 alloy. It is considered that these enriched particles correspond to a very early stage of the iron segregation found in annealed specimens which aould have taken place during cooling or during testing at 600’ C. The beryllium, however, was still fairly brittle at 600’ C when in this condition. ~icroanal~is of commercial-p~ity bery~ium (- 0.02 wt y0 iron) heat treated at 6.W C for 100 hours showed small randomly distributed iron-rich areas containing 0.08 wt y0 iron whioh were not revealed metallographioally; no segregation of iron to the grain boundaries was detected. These observations indicated that this specimen was also in an intermediate condition. 3.1.1.4.
AND
DILUTE
BERYLLIUM
ALLOYS
117
general, a close similarity between these results and those obtained on the annealed dilute-iron alloys (3.1.1.1) was evident. The 5 wt y0 iron alloy, also heat treated at 800’ C for 168 hours, was found to contain several per cent of an iron-rich phase present at the grain boundaries and within the grains (fig. 3). The large precipitates contained N 31 wt o/o iron indicating that this is the equilibrium 5 (BeisFe) phase (32.4 wt y0 iron). In addition, a fine iron-rich precipitate, lying parallel to (0001) planes, was observed within the grains. The iron content of the matrix, measured well away from any observable precipita~, was 1.1 wt %.
Fig. 3. Beryllium - 5.0 wt y0 iron alloy, heat-treated cat 800” C for 168 hours and oil qwnched, showing iron-rich precipitates. 550 x
Arc-melted 1 wt % and 5 wt % iron alloys
Preliminary results on heat-treated specimens of the 1 wt o/oiron alloy indicated that after 168 hours at 800’ C the major part of the iron is in solid solution, about 0.9 wt o/o being measured in the matrix. Iron se~egation had taken place to rando~y distribu~d cavities which were produced during the manufacture of the alloy, and to the grain boundaries. About 2 wt y. iron was measured around the cavities while the general grain-boundary enrichment was about 1.4 wt %. Up to 8 wt O$J, iron was recorded on inclusions present at the grain boundaries, and sometimes small amounts of aluminium (- 4 wt %) were also detected. In
3.1.2.1.
Beryllium-O.4 wt y. aluminium alloy
In the extruded specimen, particles 1-5 pm in diameter containing up to 80 wt o/oaluminium were present mainly at the grain’ boundaries. Thermal treatment at 700° C for 212 hours did not appear to affect the distribution of the aluminium, although iron and titanium were then found to be associated with some of the aluminium-rich precipitates. Aluminium was not detected in the matrix of either of the two specimens, indioating that the solubility of alumGum in beryllium at 700’ C is below 0.05 wt %.
118
3.1.2.2.
G.
V. T.
RANZETTA
Beryllium-O.4 wt y0 silicon alloy
AND
V.
D.
SCOTT
3.2. ELECTR~KMICROSUOPY
Most of the silioon in the as-extruded alloy Ann%aled and homogenised spscimens wer% was present at the grain bo~nd~ies in the form prepared from a beryllium-O. 15 wt y. iron alloy of precipita~s of several microns diameter, the by heat-treating the metal at 800’ C from 2 to matrix containing - 0.1 wt %. Up to 80 wt */* 120 hours and at 920’ C for 2 hours respectively, silicon was measured in typical particles, whilst and oil-quenc~ng. some awn-boundary pr~ipita~ also contained aluminium and, ocoasionally, iron and titanium. The grain boundary precipitates found in After heat treatment at 700’ C for 212 hours, th% silicon-r~oh particles ap~ared to be more these specimens are illustrated in the electron ~crograph (fig. 4) which was obta~ed by rounded in shape, and contained more iron. rep~eating a m%tallo~aphie~lly-polished and 3.123. Beryllium-O.2 wt y. c~omium alloy etched surface similar to that shown in fig. la. The results ava~able for this alloy are limited to a specimen heat-trea~d for 48 hours at 700° C and possessing good ductility when tested at 600’ C. numerous c~o~urn-rich precipitates were found, mainly at the grain bo~da~es. There appeared to be two types present, a few large (- 8 pm) precipitates tog%ther with many more small (- I pm) panicles. 30 wt o/0 ~~omium together with small amounts of iron, alurni~~ and silicon was measured on the large particles, a typical analysis giving values of 0.2 wt o/0iron, 6 wt y. aln~nium and 4 wt y. silicon. The size of the small inclusions prevented accurate analysis but N 20 wt */* chromium and N 0.15 wt o/0iron were measured while no aluminium or ailioonwas found. No general enrichment of the grain boundary in chromium could be detected. The matrix was found to contain less than 0.02 wt y. of either chromium or iron. 3.1.2.4.
Beryllium-O. I8 wt y0 manganese alloy
A sample, heat-treated for 2 hours at 800” C to give good ~gh-tem~rat~e prop%rties, was found to contain manganese-rich inctlusions present mainly at gram boundaries. 30 wt y. manganese was measured in the larger inolusions, which usually also contained alumiuium and silicon (- 4 wt */*) and occasionally iron (0.3 wt %). No general enriohment of the grain boundaries in manganese was evident. The manganese content of the matrix was 0.1 wt %, while 0.02 wt % iron was also found to be present,
Fig. 4. Beq4liu.m - 0.X wt y0 iron alloy, annealed at 800’ C for 120 hours; etched in 5 o/0RF. in glycerol. (Re@ca electron micrograph). 20000 x
No evidence of precipitation within the matrix was obtained. The fracture mod% after tensile testing at 600” C was found to depend upon annealing time and strain rate ; for example, after annealing for 2 hours at 800” C, the fracture (fig. 5) was tran~anular at a stram rate of 2.75 y. per min,
METALLOORAPHIC
STRUCTURE
OF
BERYLLIUM
AND
DILUTE
BERYLLIUM
ALLOYS
119
of some, but not all the metal grains, together with a denuded zone N 2 pm wide adjaoent to the grain boundaries (fig. 7b). From the study of a number of specimens given heat treatments from 2 to 120 hours, an impression was gained that the grain-boundary precipitates
were
larger
after
the
longer
anneal. A common feature of the microstructure was a pattern of long and nearly parallel dislocations (fig. 8a) lying in the plane of the foil, usually a bmal plane, with little evidence of interaction. Some disloctltion intersections (e.g. fig. 8b) can be seen in other regions of the specimen; at some intersections, the fourfold nodes hsve split into three-fold nodes but little dissociation of the dislocations has occurred, indicating that the stacking-fault energy of the material is fairly high.
Fig. 5.
Beryllium
-
0.16 wt y0 iron alloy, annealed
at 800’ C for 2 hours; transgranular fracture at 600” C. (Replica
electron micrograph)
5000 x
but was intergranular (fig. 6) at & strain rate of 0.2 o/o per min. Longer annealing times, however, increased the proportion of transgranular fracture at all strain rates. The intergranular fracture, (fig. 6), showed numerous precipitates (- 0.5 ,um diam) together with a slight dimpling of the surface ; a, fine background structure is also visible. A thin foil, specimen obtained from an alloy annealed at 800” C for 2 hours shows similar precipitates 0.1 ,um to 0.5 ,um thick which are seen to be plate-like, outlining the grain boundaries (fig. 7a). Diffraction patterns from the grain-boundary regions geve only Kikuchi lines arising from the neighbouring grains, the beryllium foils obtained to d&e being too thick to permit the determination of the crystal structure of these particles. Smaller precipitates, about 0.1 ,um in diameter and of density N 1014/cm3, were observed within the matrix
Fig. 6.
Beryllium
-
0.15 wt y0 iron alloy, annealed
as specified in fig. 5, but deformed at a slower strain rate; intergranular fracture at 600’ C. (Replica electron micrograph) 10 000 x
120
Q. V. T. ~A~ZETTA
la. precipitates a$ gmin boundaries. 10 000 x Fig. 7.
3.2.2.
AND V. D, SCOTT
b.
precipitates at gr&n bound&es together with smaller precipitates in matrix. 10000 x
Beryllium - 0.16 wt yc iron alloy, am&led at 800” C for 8 hours. (!lMn foil electron micrographs)
Homogenised Specimen8
fine grain-boundary precipitates present in metal given 8 prior a~~~ treatment. Inclusions were also visible at the grain A marked difference in the dislocation strucbound&ties in the alloy quenched from 920° C, Bnd replicas revealed no microstructural dif- ture was also apprtrent. It may be seen that ference between the homoge~~d and the many ~sloc~tions in the electron ~cro~~ph, (fig. 10a) are situated in piled-up groups, which annertlsd metal. indicates tha,t cross-slip processes are diff< Fig. 9 is a typical fractured GYM from a, spectimen tested at 600’ C. This intergranular and tha.tthe stacking-fault energy of the material mode of failure corresponds to low ductility and is low. The short dislocation lines lie on planes the surface is noti~&bly different from the which are inclined to the foil surface, (pool), fractured grain-boundary surface of the annealed and which intersect along a (IIPO) dire&ion ; material. “Dimpling” is much more marked the pile-up plane is a pyramidal type, most ttnd only a few large precipit&~s (sever&l probably (1Oii). An adjacent region (fig. lob) of the same foil shows an extensive hexagonal pm diameter) are visible. Precipitates like those found in the annealed network produced by screw dislocations interspecimen were not observed in thin foils pro- secting on the basaltplane. Three-fold nodes are duced from homogenised material. Furthermore, alternately extended and contracted, in contrast homoge~sation o&usedthe ~s&p~~~n~e of the to the wholly contr&c~ nodes illustrated in
METALLO~~APEIG
a.
STRUCTURE
OF
BERYLLIUM
thin foil showing disbctationslying nearly parallel to the plane of the foil. 100 000 x Fig. 8.
b.
BERYLLIUM
ALLOYS
121
intersecting dislocations with little interaction. 60 000 x
Beryllium - 0.15 wt y’ iron alloy, anneeled et 800” C for 120 hours. (Thin foil electron micrographs)
fig. 9b. A range of node radii were measured, and from these the stacking-fault energy of the material, y, was calcubted using the relation y =$@/2B, as des&bed by Whelan 2). Values obtained ranged between 18 and 50 ergs/ems, stress variations witbin the network probably accounting for these differences. 4.
ABND DILUTE
Discussion
Electron-probe microanalysis and electron microscopy have been used in the present study to investigate the distribution of impu~ties and alloying additions in beryllium and the effect that certain thermal treatments may have upon this dissemination. The experimental results, which concern particularly slloys of beryllium with iron, are considered in the first section of the discussion and related to the ductility of
the metal at 60.0’ C. In the second section of the discussion, theories which have been proposed to explain the low ductility of beryllium at high temper~t~es are critically reviewed in ctonjunction with the present results. 4.1.
4.1.1.
EXPER~E~TAL
FIERCE
~er~l~~~rn-~~on Ai%ys
Commercial-purity beryllium and berylliumiron alloys, containing up to 0.4 wt o/0 iron, exhibited low ductility a;t 600’ C in the asreaeived or homoge~sed condition but greatly improved ductility in the annealed state. A clear distinction w&s found by means of microanalysis between samples examined in the two conditions. The grain bounda~es in specimens annealed to give good ductility were found
Gt.V.
122
Fig. 9.
Beryllium
RANZETTA
- 0.16 wt y0 iron alloy, homogenised
at 920’ C for 2 hours and quenched, granular (Replica
T.
fracture
electron
showing
inter-
at 600’ C.
micrograph)
5000 x
to be enriched with iron while, in contrast, homogenised or as-extruded specimens with associated low ductility showed no grainboundary enrichment. Also inclusions containing aluminium and silicon, situated mainly at the grain boundaries of the metal, showed after annealing iron-enrichment up to a measured value of N 7 wt %, with indications that the iron had reacted with the aluminium. The amount of iron in the matrix did not change during heat treatment (experimental sensitivity of the measurements N 0.02wt %) showing that the total remount of iron involved in the segregation was small. It is thus seen that during the annealing treatment some iron had diffused to the grain boundaries, precipitates and other imperfections (for example the cavities in the 1 wt o/o iron alloy) and that this segregation of iron was
AND
V.
D.
SCOTT
accompanied by an improvement in hightemperature ductility. A change from intergranular to transgranular fracture of the specimen waz also observed to be associated with the improvement. This fracture mode was dependent upon both annealing time and strain rate, the minimum annealing time to produce higher ductility at 600’ C being found to depend on the rate at which the specimen w&s strained during tensile tests. At a strain rate of 0.2 y. per minute the ductility improvement coincided with the &at detectable iron-enrichment while at faster strain rates (e.g. 2 o/o per minute) similar ductility values were measured on specimens annealed for shorter times. Although no iron-enrichment was detected at the boundaries in the l&er specimens it was presumably present, but to a lesser extent. The presence of fine iron-rich precipitates at the grain boundaries of ductile. material, such as the particles seen by electron microscopy in the 0.15wt y. iron alloy, would not have been resolved separately in the microanalyser and, if iron-rich, could account for the continuous outlining observed. Alternatively, the observed iron-enrichment could be a form of equilibrium segregation without grain-boundary precipitation 314).It was shown that the enrichment was reversible and could be removed by thermal treatment at a temperature a little above the annealing temperature, the temperature at which this reversion first occurs varying with the iron content of the alloy 1). Since the fine-scale grain-boundary precipitates seen by electron microscopy in the 0.15 wt y. iron alloy were also removed by this heat treatment, it would appear that the iron enrichment is due to the precipitates rather than equilibrium segregation. Transmission electron microscopy of the 0.15 wt o/o iron alloy quenched from 920’ C has not revealed fine-scale precipitates or zones within the matrix of the “age-hardening type” found notably in aluminium alloys, and which might have accounted for the low ductility of the beryllium alloy at 600” C. This negative
~ETALLO~RAPHIC
a.
STRUCTURE
OF
BERYLLIUM
Dislocations formed into piled-up groups. 40 000 x
AND
b.
DILUTE
BERYLLIUM
ALLOYS
123
HexctgonaJnetwork showing extended and con. tracted nodes. 80 000 x
Fig. IO. Beryllium - 0.15 wt % iron alloy, homogenis~ at 920” C for 2 hours randquenched. (Thin foil electron micrographs)
evidence is not due to unfavourable contrast conditions sinoe experiments with a goniometer &age also failed to reveal any precipitates. E’urthermore, if an ageing reaction were to ocour, an overaged precipitate might have been expected in the annealed material but no such structure was observed. (Since the matrix precipitates seen in electron ~cro~aphe (e.g. fig. 7b) were not found in other regions of the same specimen (e.g. fig. 7a) it is unlikely that this structure is representative of the metal). Materialannealed at lower temperatures (450%) was also examined but to date no evidence 2f solid solution decomposition has been ob;ained. Microanalysis results obtained on the 1 wt y. tnd 5 wt o/0 iron alloys showed that iron was foluble in beryllium up to about 1 wt o/o at 300’ C and that any iron in excess of this amount readily p~cipita~d as the c phase 5)
containing about 32 wt % iron. Although it is very unlikely that the precipitation of c phase would take place in a homogenous be~~urn alloy of low iron oontent (e.g. below 0.4 wt %) at 800’ C, this oould perhaps ocour in the annealed specimens if sufficient segregation of free iron oocurred to cause the solubility limit for this element to be exceeded locally for example at grain boundaries. The highest local iron enrichment measured in alloys containing 0.4 wt o/o iron or less was 7 wt Oh,but as the small size of the precipitates precludes accurate analysis, these results do not rule out the possible presence of the c phase. 4.1.2.
Other Beryllium
Alloys
The ~st~bution of mangan~ in the anneaIed Be-Mn alloy was similar to that found in the iron alloys, except that no general enrichment
124
o. V. T. RANZETTA
of the grain boundaries was observed. No chromium was deteoted in the matrix of the bery~um-choke alloy, the element having precipitated in the form of numerous CrBeis particles oontaiuing a trace of iron. In general, the elements al~~ium and silicon in th8 respective alloys were present as precipitates, mainly at grain boundaries, both before and after heat treatment. After armealiug, which slightly improved high-temperature ductility, some of the precipitates were enriched to a small degree in iron, prestint as au impurity in these alloys. 4.2.
GENERALDISCUSSION
Two characteristic effeots associated with the improved ductility found at 600’ C in the annealed spwimens have be8n observed:
0) The removal to a very low level of certain impurities from the matrix, as typified by the be~~urn-0~0~~ alloy.
(ii) The general segregation of certain elements, as exemplified by iron in the beryllium-iron alloys, to fain-bound~y and other sites, with the retention in solution of a siguificant amount of the element. The attainment of the bett8r mechanical properties in specimens ~h~&c~rised by group (i) may be simply explained by assuming that an increase in matrix ductility has been produced by the p~fi~ation process. The proposal of a more complicated mechanism is required iu order to explain the observed segregation behaviour of specimens in group (ii) in relation to their mechanical properties. It is probable that both processes occur in some alloys. For instance, since the oommeroialpurity beryllium used as a basis for the alloys oontain8d impurities, some of which, notably iron, are known to embrittle beryllium, it is possible that the results on some materials (e.g. the bery~ium-s~ieon alloy) were at least partly affected by the presence of such impurities. With reference to ctommercial beryllium and beryllium-iron alloys, in which segregation of
AND
V.
D.
SCOTT
iron to graiu boundaries togetherwitha relatively high level of iron retained in the matrix was observed, several theories have be8n proposed to explain their low ductility at 600’ C and subsequent improvement by annealing, but no complexly satisfao~ry explanation appears to have been presented. The most commonly quoted mechanism is pre~pitation hardening. It has been postulatedi) that the as-extruded or homogenised metal is normally in the age-hardened condition at 600” C but that on annealing an overaged precipitate is formed which leads to softening of the metal. The present work shows that if such a reaction occurs it is not a simple ageing process associated with a binary beryllium-iron syst8m, as iron should be fully in solution at the annealiug temperature. It might therefore be proposed that other elements, for example alumni, silicon or carbon, are involved in the reaction, either precipitating independently or as compounds with iron. Th8 increased ease with which the higher iron content alloys can be annealed may be explained by assuming that the increased iron concentration either reduoes the solubility of other embattling elements in beryllium or that it aids the nucleation and growth of iron-rich precipitates. However, no evidence supporting the preaipitation theory has been obtained by electron microscopy and it would appear, therefore, that other mechanist n88d to be considered. One explanation 6) has attributed the bettei duotility of annealed beryllium to the removal of iron, a strong solid-solution hardener, from the crystal lattice by alumiuium present as a liquid-phase impurity present at the grair boundaries during annealing. Indeed, the beryl bum-chromium alloy results suggest that suck a purification mechanism could be effective as this alloy was found to possess good ductility at 600’ C in conjunction with a matrix con taiuing very little iron or chromium, However accepting the possibility that very pure beryl lium is relatively ductile at these high temper. atrues, the observed ductility of aunealed alloys of higher iron content, where iron was stil
METALLOGRAPHIC
found
in the matrix,
STRUCTURE
necessitates
OF
BERYLLIUM
an alterna-
tive explanation. Furthermore, in the present work, no reduction in the iron content of the matrix was detected in a specimen of commeroial-purity beryllium (- 0.02 wt o/o iron) which was given a heat-treatment identical to that used in the above study 6) (100 h at 650’ C), although the lower aluminium content of the samples used in the present work may account for this result. The preceding theory could be modified by supposing that the iron is removed only from regions immediately adjacent to the grain boundaries which would not have been resolved by the electron-probe. Ductile regions would therefore be produced to absorb stresses at the grain boundary which would otherwise lead to premature intergranular failure. It appears unlikely, however, that this condition would remain stable indefinitely during heat treatment as the aluminium present should eventually become saturated withiron, especially in the higher iron alloys and so be unable to absorb any more. Consequently, a denuded zone, if present at the grain boundary, would have to be formed not at the annealing temperature but either during the subsequent cooling of the specimen, or at 600” C while tensile testing, the annealing treatment being required to produce suitable nuclei for the reaction. This latter mechanism is, however, not supported by experimental evidence : holding a beryllium specimen at temperature prior to testing did not improve the ductility of an as-extruded specimen, although some slight iron segregation was found, suggesting that suitable nuclei were present. A further explanation 7) has attributed the effect of thermal treatment on ductility to the strengthening of the grain boundaries alone by the keying action of a precipitate formed during annealing, but this presupposes that the boundaries were relatively weak before the heat treatment, which is not borne out by mechanical test results (the UTS being usually higher in as-extruded material). A similar difficulty is experienced with the hypothesis that the intergranular failure of
AND
DILUTE
as-extruded
BERYLLIUM
ALLOYS
125
metal is caused by the weakening
effect of a soft or molten constituent example aluminium or aluminium-silicon)
(for at
the grain boundaries. Further, although such molten phases have been found in certain beryllium alloys 8) and may be assumed to combine with iron and become more stable upon annealing, the hypothesis is not consistent with the effect upon beryllium of small aluminium additions. These did not result in a highly brittle alloy, although in accordance with this theory more molten material should have been present at 600’ C. It may be seen, therefore, that the various mechanisms which have been proposed to explain the high-temperature ductility of beryllium may be broadly classified into two main groups, the first concerned with the matrix properties of the metal and the second with its grain-boundary properties. However, no one simple theory appears to adequately explain all the experimental evidence, and it is therefore possible that both matrix and grain-boundary mechanisms are operative. Although it may be argued that the iron-segregation phenomenon with which the mechanisms discussed above are mainly concerned is only incidental to the change in mechanical properties, this is unlikely since ductility and fracture mode both appear to be closely related to this segregation 9). With reference to matrix ductility, apart from the beryllium-chromium alloys little change in the matrix structure which could account for the observed matrix softening 9) was found upon annealing. Certainly no metallographic evidence is currently available to support a precipitation-ageing hypothesis and the softening is therefore more consistent with the removal of embrittling elements from solution to the grain boundaries. There is some evidence which suggests that solid-solution hardening may occur in the case of beryllium alloys of higher iron content. Thus the differing degrees of dislocation interaction revealed in the electron micrographs (figs. 8 and 10) is indicative of different values for the stacking-fault energy of the 0.15 wt o/o alloy
126
o.
V.
T.
RANZETTA
in the annealed and homogenised conditions ; for example, quenching from 800" C resulted in higher values than were found after quenching this alloy from 920’ C. To some degree the differences are in accord with the mechanical properties of the two materials, since the latter condition was found by tensile testing to be more brittle in the temperature range 300” C to 600” C. This solid solution hardening may also be present in commercial-purity beryllium but to a lesser extent. Considering the grain boundaries - if these are brittle in homogenised metal, as both the mode of tensile fracture and the high tensile strength suggest, it is possible that they are strengthened on annealing. This could perhaps result from the removal of an embrittling thin film or ultra fine precipitates present at the boundary and the consequent formation of a coarser precipitate such as that observed in thin foils of annealed material. Should these precipitates act as dislocation sources, the ease with which dislocations sre generated could depend critically upon the precipitate distribution and the process may be more readily operative in annealed specimens. Consequently, in this material the propagation of plastic flow into neighbouring grains could be facilitated, with a corresponding relaxation of stress at the grain boundaries, and this would lead to the higher ductility and reduced tendency for intergranular failure which are observed experimentally. In conclusion, considerable experimental evidence is now being accumulated concerning the mechanical behaviour of beryllium and its related microstructure and the effect that thermal treatment may have upon this, but a comprehensive solution to the ductility problem in beryllium has not yet been achieved. For example, it is still uncertain whether or not very high purity beryllium would be ductile, although the present results on berylliumchromium alloys suggest that some improve-
AND
V.
D.
SCOTT
ment of ductility in the high-temperature range ms,y be achieved by purification of the matrix. It would be particularly desirable to extend the present work to investigate the effeot that carbon, oxygen and nitrogen may have upon the properties of beryllium, and also to include some studies on ternary systems. Future work will be espeoislly with the properties of dislocations
concerned and their
distribution as a function both of temperature and of metal purity, and will be extended to lower temperatures, which are of particular interest to the metal technologist. Acknowledgements Acknowledgements sre due to the staff of the beryllium section under Dr. A. J. Martin, who provided the material and mechanical-test data and carried out the necessary heat treatments and specimen preparation. Thanks are also due to Professor J. Nutting, Dr. A. Moore and Mr. F. Morrow for helpful discussions. This work was carried out in the Research Metallurgy Branch, AWRE Aldermaston, under the direction of Mr. G. C. Ellis. References
1)
A. Moore, F. Morrow, V. D. Scott and D. A. Cheer, Conf. on Metallurgy London,
of Beryllium
(Inst. of Metals,
1961)
M. J. Whelan,
Proc. Roy
Sot.
A249 (1958) 114
Y 3)
D. McLean,
4,
C. Crussard, J. Plateau and G. Henry,
6)
Saclay, 1960) R. J. Teitel and
9
Univ.
Grain Boundaries
Press,
Y.
Adda,
Maurice
N.
(Oxford
1957) p. 116
of Grain Boundaries
Mining, Met.,
in Metals
(4th Metallurgy
M.
Petrol., Azam,
Cohen, Engrs. G.
Properties Colloquium,
Trans.
Am.
185 (1949)
Donze,
J.
Inst.
285
Mallen,
and M. Weisz, Compt. Rend.
Acad.
F. Sci.
‘1
(Paris), 254 (1962) 1052 A. J. Martin and G. C. Ellis, Conf. on Metallurgy
9
G. Donze,
9)
J. Mat. Nucl. 6 (1962) 137 A. Brown, F. Morrow and A. J. Martin, J. Less-
of Beryllium
Common
(Inst.
of Metals,
D. Dutilloy,
Metals
London,
R. le Hazif
3 (1961)
62
1961)
et Y.
Adda,
THE
OF NUCLEAR
EFFECT
~~
lo,2
OF TEENY
(1963)
TREA~ENT
BER~LI~
~OR~-EO~~D
113426,
ON THE ~ETALLOGRAPHIC
AND DILUTE
BERYLLIUM
C. V. T. RANZETTA Metallurgy
.lXv&iorc, Atomic
Wmpons
PUBLIS~G
CO.,
AMSTE~~
STRUCTURE
OF
ALLOYS
and V. D. SCOTT
Research Establiahmmt,
Akkrmmtm,
Be&a.,
UK
Received 13 March 1963
Electron microscopy and electron-probemicroanalysis techniqueshave been used to study t,hemiorostructure of beryllium and some dilute beryllium alloys. The effect of thermal treatment on the dissemination of impuritiesand alloying elements has been investigated with particular referenceto the mechanical properties of the metal at 600” C. In beryllium-iron alloys annealed to improve hightemperature ductility the structure was typified by the presence of large iron-containing grain-boundary precipitates, together with a more general segregation of iron to grain boundariesand other favourable sites. Some indication of solid-solution hardening in this system was apparentbut no evidenceof a precipitationageing process was found. It is therefore suggested that the microstruct.urein the grain-boundary regions of the metal may be important in affecting the ductility and coherency of metal grains at high temperatures. Annealed beryllium-ohromium alloys had good high-temperatureductility which in this case may be associated also with the observed purification of the matrix.
La microscopic Blectronique et la technique de micro-analyse par sonde Blectroniqueont Bte utilisdes pour Studier la microstructure du b&yllium et de quelques alliages peu oharg& en Elements d’addition de beryllium. L’influenoe du traitement thermique sur la repartition des impuretes et des Elements d’alliage a Bte t%udit%een liaison Btroite avec les propri&tb mecaniques du m&al it 600” C. Dans les alliages b&yllium-fer, recuits dans le but d’ameliorer la duotilite it haute temperature, la structure Btait caract&isee par la presenoe de gros precipites riches en fer aux joints de grains ainsi que par une segregation plus g&&ale du fer, soit aux joints de grains, soit au hasard. On a mis en 6vidence
un certain durcissement structural dans cet alliage, mais on n’a pas trouve de preuve caraoteristiquedu processus de vieillissement aocompagne de precipitations. C’est pourquoi on a suggere que la, ductilitb du metal et la cohesion entre grains it temperature &e&e devraient d$pendre $troitement de la mierostructure existant au voisinage des contours de grams. Les alliages Be-Cr recuits possedent une bonne ductilite a haute temp&ature, fait qui peut Gtre Bgalementrelit5&la purificationobserveede la matrice.
Es wurden elektrone~~oskopische und elektronenm~o~al~~che Verfahren zur ~ntersueh~g des Gefiigesvon Berylliumund einigenl&lichen BerylliumLegierungen benutzt. Der Einfluss der Warmebehandlung auf die Verteilung von Verunreinigungen und Legierungselementen wurde untersucht, wobei besonders auf die mechanisohen Eigenschaften des Metalls bei 600” C geachtet wurde. In Beryllium-Eisen-Legierungen, welche zwecks Verbessernngihrer Hoch~mperaturdu~ilit~t warmebehandelt wurden, war die Struktur typisiert durch die Gegenwartvon grosseneisenhaltigenKorngrenzenniederschlagenund einer mehr allgemeinen Abscheidung von Eisen an den Korngrenzen und anderen Stellen. Einige Anzeichen von Mischkristallh&tung in diesem System wurden festgestellt. Es konnte jedoch keine Klarheit iiber einen Niederschlagalterungsprozess gefunden werden. Es wurde daher ~geno~en, dass das Gefiige in den Korngrenzemegionen von wesentlichem Einfiuss auf die Dehnbarkeit und Koharenz der Metallkorngrenzen bei hohen Temperaturen ist. Die warmebehandelten Beryllium-Chrom-Legierungen hatten gute Hochtemperaturduktilitiit, was in diesem Fall such zusammenbangen mag mit der beobachteten Reinheit der Matrix.
113
114 1.
G. V. T. RANZETTA AND V. D. SCOTT Introduction
A considerable amount of evidence has been accumulated during the past few years showing that the mechanical properties of beryllium in the temperature range 400-600” C depend strongly on the composition and thermal history of the metal. It may be said generally that rapid cooling of commercial-purity beryllium or dilute beryllium alloys from high temperatures, such as those used for extrusion (- 1050’ C), leads to a fall in ductility with increase in temperature above 400’ C, to a ~~rnurn at about 600” C. This drop in ductility can be avoided by a slow cooling or annealing treatment which is usually more effective when applied to dilute beryllium alloys rather than to commercial-purity beryllium l). In the present paper, results obtained by the application of electron-probe microanalysis and electron microscopy to the study of such effects in beryllium and beryllium alloys are presented. Since tensile test data were available for most of the specimens examined it has been possible to correlate microstructural information with mechanical behaviour and to discuss the results in relation to previous theories proposed to explain this ductility phenomenon.
2. 2.1.
Experimental Details ORHXN OF MATERIAL
The bery~i~ used in this study was produced from Pechiney commercial-purity electrolytic flake which was melted, together with any desired additions, cast, and subsequently hotextruded to a size suitable for a mechanical evaluation of its properties at 600* C. In addition a range of binary and ternary alloys was made by arc-melting together small quantities of powdered flake beryllium and the required alloying element. This group of specimens was intended solely for microanalytical studies and was in a form u~uitable for mechanical testing.
ELECTRON-PROBE MICROANALYSIS
2.2.
Microanalysis was carried out using a Cambridge scanning electron probe inst~ment, quantitative measurements being made in most cases by reference to pure metal standards. 2.3.
ELECTRONBIICROSCOPY
Replica methods have been used to study the surface of both metallographioally-prepared and fractured specimens of a beryllium0.15 wt y0 iron alloy. In addition, thin foils of this metal have been prepared by chemical and electroche~cal dissolution for direct examination in the electron microscope. 3.
Results
To avoid the unnecessa~ repetition 03 thermal treatment details, specimens heat treated to give better ductility at 600° C will be referred to as “annealed” and those quenched from high temperatures, giving a less ductile product, as “homogenised”. 3.1.
ELECTRON-PROBE MICROANALYSIS
In general the low level of impurities in commercial-purity beryllium and beryllium-iron alloys has made the study of their dissemination ~fficult and dilute alloys consisting of bery~um and certain of the commonly-occnrring impurity elements have therefore been examined. With regard to quantitative microanalysis the volume of material irradiated by the electron probe will be about 30 prna depending upon analysis conditions. As a consequence the probe values quoted for inclusions smaller than about 4 pm in diameter may tend to be low; this size effect could be considerable for particles with a diameter of less than 2 ,um. X1.1.
Beryrllium- Iron Alloys
Alloys containing iron in the range 0.4 wt y. to a nominal 0.03 wt %, as found in commercial grade beryllium, have been investigated in both the “annealed” (ductile) and “homogeni~’ (brittle) condition.
~~TALLO~RA~H~C
STRUCTURE
OF
BERYLLIUM
s.
Fig. 1.
b.
Electron.
d.
X-my,
Ax.
1400 x
1400 x
Optical.
AND
DILUTE
BERYLLIUM
ALLOYS
400 X
0.
X-my,
Fe.
1400 X
8.
X-my,
Si.
1400 X
Beryllium - 0.18 wt % iron alloy, annealed at 800’ G for 120 houra and slowly cooled, &owing segregation ati gmin boundaries.
1x6
116
G.
a.
Fig. 2.
Electron.
Beryllium
V.
T.
RANZETTA
‘1000 x -
AnneaM
V.
D.
SCOTT
b.
X-ray,
Al.
0.15 wt y0 iron alloy, homogenised at 1000” C for quenched, showing alumiuium at grain boundaries.
Some preliminrtry results obtdned on alloys containing greater amounts of iron (1 wt */*and 5 wt %) sre also presented, 3.1.1.1.
AND
specimens (Ductile at
600” C)
A general enrichment of the grain boundaries in iron was found in these specimens. This was most pronounced in the higher iron alloys, 0.8 wt y. Fe being indicated at the grain boundary of a 0.4 wt y. alloy, but was still just deteotable in a 0.08 wt y. specimen. The grain boundary segregation of iron is clearly shown in figs. 1 a-c, which were obtained from a 0.18 wt y. iron alloy that had been annealed at 600’ C for 120 hours and slowly cooled. The degree of iron enrichment increased with time during heat-treatment, indicating that iron is slowly diffusing to the grain boundaries during annealing. For example, in the case of a 0.16 wt o/0 iron alloy heat-treated at 800” C, no iron was found after 2 hours, slight enrichment w&s visible after 10 hours, and after 100 hours the iron WM readily detectable. P&icles 1 to 3 ,.umin size and located mainly at grain bound~ies were fairly common in annealed beryllium specimens, ~ero~n~lysis showing the presence of iron in concentrations of up to 7 wt Oh, usually together with a few per cent aluminium and silicon (figs. 1 d-e).
1000 x 2 hours
and oil
A slight enrichment of the grain boundaries in aluminium may be observed which appears to be associated with the ~stribution of iron, but not every grain boundary was enriched to the same extent. The iron con~nt of the matrix was not found to be significantly affected by the heat treatment; a very small change ( < 0.02 wt %) would be within the experimental scatter of most of the measurements and remain undetected. 3.1.1.2. Homogenised specimens (Low ductility at 600’ C) Specimens of as-extruded and annealed beryllium which h&d been subsequently homoge~sed at high tem~r&t~es (90~1060’ C) and quenched &ve been examined. No grain-boundary enrichment was detected, the iron being distributed uniformly throughout the specimen irrespeotive of its initial dissemination; thus the iron segregation at the grain boundaries associated with increased ductility a;t 600’ C is reversible. Figs. 2a, and 2b were taken from it, grain boundary showing a, rela;tively high concentration of alu~~um-rich inclusions in a commerci&l-purity beryllium specimen that was oil-quenched after heat treatment at 1000” C. Up to 60 wt y0 aluminium was measured in these inclusions, no other elements being detected.
METALLOGRAPHIC
STRUCTURE
OF
BERYLLIUM
3.1.1.3. Specimens in an intermediate condition During extrusion and cooling, beryllium specimens are subjected to a form of heat treatment that ~a~ot be closely defined and it might therefore be expected that as-extruded specimens oould be in a condition in~~e~ate between the homogenised and annealed state. However, microanalysis results from such alloys were similar to those obtained for homogenised metal. No segregation of iron to the grain boundaries was observed but it was possible to locate inclusions slightly richer in iron than the matrix ; for example 0.3 wt $.&Fe was measured in an inclusion contained in an 0.18 wt y0 alloy. It is considered that these enriched particles correspond to a very early stage of the iron segregation found in annealed specimens which aould have taken place during cooling or during testing at 600’ C. The beryllium, however, was still fairly brittle at 600’ C when in this condition. ~icroanal~is of commercial-p~ity bery~ium (- 0.02 wt y0 iron) heat treated at 6.W C for 100 hours showed small randomly distributed iron-rich areas containing 0.08 wt y0 iron whioh were not revealed metallographioally; no segregation of iron to the grain boundaries was detected. These observations indicated that this specimen was also in an intermediate condition. 3.1.1.4.
AND
DILUTE
BERYLLIUM
ALLOYS
117
general, a close similarity between these results and those obtained on the annealed dilute-iron alloys (3.1.1.1) was evident. The 5 wt y0 iron alloy, also heat treated at 800’ C for 168 hours, was found to contain several per cent of an iron-rich phase present at the grain boundaries and within the grains (fig. 3). The large precipitates contained N 31 wt o/o iron indicating that this is the equilibrium 5 (BeisFe) phase (32.4 wt y0 iron). In addition, a fine iron-rich precipitate, lying parallel to (0001) planes, was observed within the grains. The iron content of the matrix, measured well away from any observable precipita~, was 1.1 wt %.
Fig. 3. Beryllium - 5.0 wt y0 iron alloy, heat-treated cat 800” C for 168 hours and oil qwnched, showing iron-rich precipitates. 550 x
Arc-melted 1 wt % and 5 wt % iron alloys
Preliminary results on heat-treated specimens of the 1 wt o/oiron alloy indicated that after 168 hours at 800’ C the major part of the iron is in solid solution, about 0.9 wt o/o being measured in the matrix. Iron se~egation had taken place to rando~y distribu~d cavities which were produced during the manufacture of the alloy, and to the grain boundaries. About 2 wt y. iron was measured around the cavities while the general grain-boundary enrichment was about 1.4 wt %. Up to 8 wt O$J, iron was recorded on inclusions present at the grain boundaries, and sometimes small amounts of aluminium (- 4 wt %) were also detected. In
3.1.2.1.
Beryllium-O.4 wt y. aluminium alloy
In the extruded specimen, particles 1-5 pm in diameter containing up to 80 wt o/oaluminium were present mainly at the grain’ boundaries. Thermal treatment at 700° C for 212 hours did not appear to affect the distribution of the aluminium, although iron and titanium were then found to be associated with some of the aluminium-rich precipitates. Aluminium was not detected in the matrix of either of the two specimens, indioating that the solubility of alumGum in beryllium at 700’ C is below 0.05 wt %.
118
3.1.2.2.
G.
V. T.
RANZETTA
Beryllium-O.4 wt y0 silicon alloy
AND
V.
D.
SCOTT
3.2. ELECTR~KMICROSUOPY
Most of the silioon in the as-extruded alloy Ann%aled and homogenised spscimens wer% was present at the grain bo~nd~ies in the form prepared from a beryllium-O. 15 wt y. iron alloy of precipita~s of several microns diameter, the by heat-treating the metal at 800’ C from 2 to matrix containing - 0.1 wt %. Up to 80 wt */* 120 hours and at 920’ C for 2 hours respectively, silicon was measured in typical particles, whilst and oil-quenc~ng. some awn-boundary pr~ipita~ also contained aluminium and, ocoasionally, iron and titanium. The grain boundary precipitates found in After heat treatment at 700’ C for 212 hours, th% silicon-r~oh particles ap~ared to be more these specimens are illustrated in the electron ~crograph (fig. 4) which was obta~ed by rounded in shape, and contained more iron. rep~eating a m%tallo~aphie~lly-polished and 3.123. Beryllium-O.2 wt y. c~omium alloy etched surface similar to that shown in fig. la. The results ava~able for this alloy are limited to a specimen heat-trea~d for 48 hours at 700° C and possessing good ductility when tested at 600’ C. numerous c~o~urn-rich precipitates were found, mainly at the grain bo~da~es. There appeared to be two types present, a few large (- 8 pm) precipitates tog%ther with many more small (- I pm) panicles. 30 wt o/0 ~~omium together with small amounts of iron, alurni~~ and silicon was measured on the large particles, a typical analysis giving values of 0.2 wt o/0iron, 6 wt y. aln~nium and 4 wt y. silicon. The size of the small inclusions prevented accurate analysis but N 20 wt */* chromium and N 0.15 wt o/0iron were measured while no aluminium or ailioonwas found. No general enrichment of the grain boundary in chromium could be detected. The matrix was found to contain less than 0.02 wt y. of either chromium or iron. 3.1.2.4.
Beryllium-O. I8 wt y0 manganese alloy
A sample, heat-treated for 2 hours at 800” C to give good ~gh-tem~rat~e prop%rties, was found to contain manganese-rich inctlusions present mainly at gram boundaries. 30 wt y. manganese was measured in the larger inolusions, which usually also contained alumiuium and silicon (- 4 wt */*) and occasionally iron (0.3 wt %). No general enriohment of the grain boundaries in manganese was evident. The manganese content of the matrix was 0.1 wt %, while 0.02 wt % iron was also found to be present,
Fig. 4. Beq4liu.m - 0.X wt y0 iron alloy, annealed at 800’ C for 120 hours; etched in 5 o/0RF. in glycerol. (Re@ca electron micrograph). 20000 x
No evidence of precipitation within the matrix was obtained. The fracture mod% after tensile testing at 600” C was found to depend upon annealing time and strain rate ; for example, after annealing for 2 hours at 800” C, the fracture (fig. 5) was tran~anular at a stram rate of 2.75 y. per min,
METALLOORAPHIC
STRUCTURE
OF
BERYLLIUM
AND
DILUTE
BERYLLIUM
ALLOYS
119
of some, but not all the metal grains, together with a denuded zone N 2 pm wide adjaoent to the grain boundaries (fig. 7b). From the study of a number of specimens given heat treatments from 2 to 120 hours, an impression was gained that the grain-boundary precipitates
were
larger
after
the
longer
anneal. A common feature of the microstructure was a pattern of long and nearly parallel dislocations (fig. 8a) lying in the plane of the foil, usually a bmal plane, with little evidence of interaction. Some disloctltion intersections (e.g. fig. 8b) can be seen in other regions of the specimen; at some intersections, the fourfold nodes hsve split into three-fold nodes but little dissociation of the dislocations has occurred, indicating that the stacking-fault energy of the material is fairly high.
Fig. 5.
Beryllium
-
0.16 wt y0 iron alloy, annealed
at 800’ C for 2 hours; transgranular fracture at 600” C. (Replica
electron micrograph)
5000 x
but was intergranular (fig. 6) at & strain rate of 0.2 o/o per min. Longer annealing times, however, increased the proportion of transgranular fracture at all strain rates. The intergranular fracture, (fig. 6), showed numerous precipitates (- 0.5 ,um diam) together with a slight dimpling of the surface ; a, fine background structure is also visible. A thin foil, specimen obtained from an alloy annealed at 800” C for 2 hours shows similar precipitates 0.1 ,um to 0.5 ,um thick which are seen to be plate-like, outlining the grain boundaries (fig. 7a). Diffraction patterns from the grain-boundary regions geve only Kikuchi lines arising from the neighbouring grains, the beryllium foils obtained to d&e being too thick to permit the determination of the crystal structure of these particles. Smaller precipitates, about 0.1 ,um in diameter and of density N 1014/cm3, were observed within the matrix
Fig. 6.
Beryllium
-
0.15 wt y0 iron alloy, annealed
as specified in fig. 5, but deformed at a slower strain rate; intergranular fracture at 600’ C. (Replica electron micrograph) 10 000 x
120
Q. V. T. ~A~ZETTA
la. precipitates a$ gmin boundaries. 10 000 x Fig. 7.
3.2.2.
AND V. D, SCOTT
b.
precipitates at gr&n bound&es together with smaller precipitates in matrix. 10000 x
Beryllium - 0.16 wt yc iron alloy, am&led at 800” C for 8 hours. (!lMn foil electron micrographs)
Homogenised Specimen8
fine grain-boundary precipitates present in metal given 8 prior a~~~ treatment. Inclusions were also visible at the grain A marked difference in the dislocation strucbound&ties in the alloy quenched from 920° C, Bnd replicas revealed no microstructural dif- ture was also apprtrent. It may be seen that ference between the homoge~~d and the many ~sloc~tions in the electron ~cro~~ph, (fig. 10a) are situated in piled-up groups, which annertlsd metal. indicates tha,t cross-slip processes are diff< Fig. 9 is a typical fractured GYM from a, spectimen tested at 600’ C. This intergranular and tha.tthe stacking-fault energy of the material mode of failure corresponds to low ductility and is low. The short dislocation lines lie on planes the surface is noti~&bly different from the which are inclined to the foil surface, (pool), fractured grain-boundary surface of the annealed and which intersect along a (IIPO) dire&ion ; material. “Dimpling” is much more marked the pile-up plane is a pyramidal type, most ttnd only a few large precipit&~s (sever&l probably (1Oii). An adjacent region (fig. lob) of the same foil shows an extensive hexagonal pm diameter) are visible. Precipitates like those found in the annealed network produced by screw dislocations interspecimen were not observed in thin foils pro- secting on the basaltplane. Three-fold nodes are duced from homogenised material. Furthermore, alternately extended and contracted, in contrast homoge~sation o&usedthe ~s&p~~~n~e of the to the wholly contr&c~ nodes illustrated in
METALLO~~APEIG
a.
STRUCTURE
OF
BERYLLIUM
thin foil showing disbctationslying nearly parallel to the plane of the foil. 100 000 x Fig. 8.
b.
BERYLLIUM
ALLOYS
121
intersecting dislocations with little interaction. 60 000 x
Beryllium - 0.15 wt y’ iron alloy, anneeled et 800” C for 120 hours. (Thin foil electron micrographs)
fig. 9b. A range of node radii were measured, and from these the stacking-fault energy of the material, y, was calcubted using the relation y =$@/2B, as des&bed by Whelan 2). Values obtained ranged between 18 and 50 ergs/ems, stress variations witbin the network probably accounting for these differences. 4.
ABND DILUTE
Discussion
Electron-probe microanalysis and electron microscopy have been used in the present study to investigate the distribution of impu~ties and alloying additions in beryllium and the effect that certain thermal treatments may have upon this dissemination. The experimental results, which concern particularly slloys of beryllium with iron, are considered in the first section of the discussion and related to the ductility of
the metal at 60.0’ C. In the second section of the discussion, theories which have been proposed to explain the low ductility of beryllium at high temper~t~es are critically reviewed in ctonjunction with the present results. 4.1.
4.1.1.
EXPER~E~TAL
FIERCE
~er~l~~~rn-~~on Ai%ys
Commercial-purity beryllium and berylliumiron alloys, containing up to 0.4 wt o/0 iron, exhibited low ductility a;t 600’ C in the asreaeived or homoge~sed condition but greatly improved ductility in the annealed state. A clear distinction w&s found by means of microanalysis between samples examined in the two conditions. The grain bounda~es in specimens annealed to give good ductility were found
Gt.V.
122
Fig. 9.
Beryllium
RANZETTA
- 0.16 wt y0 iron alloy, homogenised
at 920’ C for 2 hours and quenched, granular (Replica
T.
fracture
electron
showing
inter-
at 600’ C.
micrograph)
5000 x
to be enriched with iron while, in contrast, homogenised or as-extruded specimens with associated low ductility showed no grainboundary enrichment. Also inclusions containing aluminium and silicon, situated mainly at the grain boundaries of the metal, showed after annealing iron-enrichment up to a measured value of N 7 wt %, with indications that the iron had reacted with the aluminium. The amount of iron in the matrix did not change during heat treatment (experimental sensitivity of the measurements N 0.02wt %) showing that the total remount of iron involved in the segregation was small. It is thus seen that during the annealing treatment some iron had diffused to the grain boundaries, precipitates and other imperfections (for example the cavities in the 1 wt o/o iron alloy) and that this segregation of iron was
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accompanied by an improvement in hightemperature ductility. A change from intergranular to transgranular fracture of the specimen waz also observed to be associated with the improvement. This fracture mode was dependent upon both annealing time and strain rate, the minimum annealing time to produce higher ductility at 600’ C being found to depend on the rate at which the specimen w&s strained during tensile tests. At a strain rate of 0.2 y. per minute the ductility improvement coincided with the &at detectable iron-enrichment while at faster strain rates (e.g. 2 o/o per minute) similar ductility values were measured on specimens annealed for shorter times. Although no iron-enrichment was detected at the boundaries in the l&er specimens it was presumably present, but to a lesser extent. The presence of fine iron-rich precipitates at the grain boundaries of ductile. material, such as the particles seen by electron microscopy in the 0.15wt y. iron alloy, would not have been resolved separately in the microanalyser and, if iron-rich, could account for the continuous outlining observed. Alternatively, the observed iron-enrichment could be a form of equilibrium segregation without grain-boundary precipitation 314).It was shown that the enrichment was reversible and could be removed by thermal treatment at a temperature a little above the annealing temperature, the temperature at which this reversion first occurs varying with the iron content of the alloy 1). Since the fine-scale grain-boundary precipitates seen by electron microscopy in the 0.15 wt y. iron alloy were also removed by this heat treatment, it would appear that the iron enrichment is due to the precipitates rather than equilibrium segregation. Transmission electron microscopy of the 0.15 wt o/o iron alloy quenched from 920’ C has not revealed fine-scale precipitates or zones within the matrix of the “age-hardening type” found notably in aluminium alloys, and which might have accounted for the low ductility of the beryllium alloy at 600” C. This negative
~ETALLO~RAPHIC
a.
STRUCTURE
OF
BERYLLIUM
Dislocations formed into piled-up groups. 40 000 x
AND
b.
DILUTE
BERYLLIUM
ALLOYS
123
HexctgonaJnetwork showing extended and con. tracted nodes. 80 000 x
Fig. IO. Beryllium - 0.15 wt % iron alloy, homogenis~ at 920” C for 2 hours randquenched. (Thin foil electron micrographs)
evidence is not due to unfavourable contrast conditions sinoe experiments with a goniometer &age also failed to reveal any precipitates. E’urthermore, if an ageing reaction were to ocour, an overaged precipitate might have been expected in the annealed material but no such structure was observed. (Since the matrix precipitates seen in electron ~cro~aphe (e.g. fig. 7b) were not found in other regions of the same specimen (e.g. fig. 7a) it is unlikely that this structure is representative of the metal). Materialannealed at lower temperatures (450%) was also examined but to date no evidence 2f solid solution decomposition has been ob;ained. Microanalysis results obtained on the 1 wt y. tnd 5 wt o/0 iron alloys showed that iron was foluble in beryllium up to about 1 wt o/o at 300’ C and that any iron in excess of this amount readily p~cipita~d as the c phase 5)
containing about 32 wt % iron. Although it is very unlikely that the precipitation of c phase would take place in a homogenous be~~urn alloy of low iron oontent (e.g. below 0.4 wt %) at 800’ C, this oould perhaps ocour in the annealed specimens if sufficient segregation of free iron oocurred to cause the solubility limit for this element to be exceeded locally for example at grain boundaries. The highest local iron enrichment measured in alloys containing 0.4 wt o/o iron or less was 7 wt Oh,but as the small size of the precipitates precludes accurate analysis, these results do not rule out the possible presence of the c phase. 4.1.2.
Other Beryllium
Alloys
The ~st~bution of mangan~ in the anneaIed Be-Mn alloy was similar to that found in the iron alloys, except that no general enrichment
124
o. V. T. RANZETTA
of the grain boundaries was observed. No chromium was deteoted in the matrix of the bery~um-choke alloy, the element having precipitated in the form of numerous CrBeis particles oontaiuing a trace of iron. In general, the elements al~~ium and silicon in th8 respective alloys were present as precipitates, mainly at grain boundaries, both before and after heat treatment. After armealiug, which slightly improved high-temperature ductility, some of the precipitates were enriched to a small degree in iron, prestint as au impurity in these alloys. 4.2.
GENERALDISCUSSION
Two characteristic effeots associated with the improved ductility found at 600’ C in the annealed spwimens have be8n observed:
0) The removal to a very low level of certain impurities from the matrix, as typified by the be~~urn-0~0~~ alloy.
(ii) The general segregation of certain elements, as exemplified by iron in the beryllium-iron alloys, to fain-bound~y and other sites, with the retention in solution of a siguificant amount of the element. The attainment of the bett8r mechanical properties in specimens ~h~&c~rised by group (i) may be simply explained by assuming that an increase in matrix ductility has been produced by the p~fi~ation process. The proposal of a more complicated mechanism is required iu order to explain the observed segregation behaviour of specimens in group (ii) in relation to their mechanical properties. It is probable that both processes occur in some alloys. For instance, since the oommeroialpurity beryllium used as a basis for the alloys oontain8d impurities, some of which, notably iron, are known to embrittle beryllium, it is possible that the results on some materials (e.g. the bery~ium-s~ieon alloy) were at least partly affected by the presence of such impurities. With reference to ctommercial beryllium and beryllium-iron alloys, in which segregation of
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iron to graiu boundaries togetherwitha relatively high level of iron retained in the matrix was observed, several theories have be8n proposed to explain their low ductility at 600’ C and subsequent improvement by annealing, but no complexly satisfao~ry explanation appears to have been presented. The most commonly quoted mechanism is pre~pitation hardening. It has been postulatedi) that the as-extruded or homogenised metal is normally in the age-hardened condition at 600” C but that on annealing an overaged precipitate is formed which leads to softening of the metal. The present work shows that if such a reaction occurs it is not a simple ageing process associated with a binary beryllium-iron syst8m, as iron should be fully in solution at the annealiug temperature. It might therefore be proposed that other elements, for example alumni, silicon or carbon, are involved in the reaction, either precipitating independently or as compounds with iron. Th8 increased ease with which the higher iron content alloys can be annealed may be explained by assuming that the increased iron concentration either reduoes the solubility of other embattling elements in beryllium or that it aids the nucleation and growth of iron-rich precipitates. However, no evidence supporting the preaipitation theory has been obtained by electron microscopy and it would appear, therefore, that other mechanist n88d to be considered. One explanation 6) has attributed the bettei duotility of annealed beryllium to the removal of iron, a strong solid-solution hardener, from the crystal lattice by alumiuium present as a liquid-phase impurity present at the grair boundaries during annealing. Indeed, the beryl bum-chromium alloy results suggest that suck a purification mechanism could be effective as this alloy was found to possess good ductility at 600’ C in conjunction with a matrix con taiuing very little iron or chromium, However accepting the possibility that very pure beryl lium is relatively ductile at these high temper. atrues, the observed ductility of aunealed alloys of higher iron content, where iron was stil
METALLOGRAPHIC
found
in the matrix,
STRUCTURE
necessitates
OF
BERYLLIUM
an alterna-
tive explanation. Furthermore, in the present work, no reduction in the iron content of the matrix was detected in a specimen of commeroial-purity beryllium (- 0.02 wt o/o iron) which was given a heat-treatment identical to that used in the above study 6) (100 h at 650’ C), although the lower aluminium content of the samples used in the present work may account for this result. The preceding theory could be modified by supposing that the iron is removed only from regions immediately adjacent to the grain boundaries which would not have been resolved by the electron-probe. Ductile regions would therefore be produced to absorb stresses at the grain boundary which would otherwise lead to premature intergranular failure. It appears unlikely, however, that this condition would remain stable indefinitely during heat treatment as the aluminium present should eventually become saturated withiron, especially in the higher iron alloys and so be unable to absorb any more. Consequently, a denuded zone, if present at the grain boundary, would have to be formed not at the annealing temperature but either during the subsequent cooling of the specimen, or at 600” C while tensile testing, the annealing treatment being required to produce suitable nuclei for the reaction. This latter mechanism is, however, not supported by experimental evidence : holding a beryllium specimen at temperature prior to testing did not improve the ductility of an as-extruded specimen, although some slight iron segregation was found, suggesting that suitable nuclei were present. A further explanation 7) has attributed the effect of thermal treatment on ductility to the strengthening of the grain boundaries alone by the keying action of a precipitate formed during annealing, but this presupposes that the boundaries were relatively weak before the heat treatment, which is not borne out by mechanical test results (the UTS being usually higher in as-extruded material). A similar difficulty is experienced with the hypothesis that the intergranular failure of
AND
DILUTE
as-extruded
BERYLLIUM
ALLOYS
125
metal is caused by the weakening
effect of a soft or molten constituent example aluminium or aluminium-silicon)
(for at
the grain boundaries. Further, although such molten phases have been found in certain beryllium alloys 8) and may be assumed to combine with iron and become more stable upon annealing, the hypothesis is not consistent with the effect upon beryllium of small aluminium additions. These did not result in a highly brittle alloy, although in accordance with this theory more molten material should have been present at 600’ C. It may be seen, therefore, that the various mechanisms which have been proposed to explain the high-temperature ductility of beryllium may be broadly classified into two main groups, the first concerned with the matrix properties of the metal and the second with its grain-boundary properties. However, no one simple theory appears to adequately explain all the experimental evidence, and it is therefore possible that both matrix and grain-boundary mechanisms are operative. Although it may be argued that the iron-segregation phenomenon with which the mechanisms discussed above are mainly concerned is only incidental to the change in mechanical properties, this is unlikely since ductility and fracture mode both appear to be closely related to this segregation 9). With reference to matrix ductility, apart from the beryllium-chromium alloys little change in the matrix structure which could account for the observed matrix softening 9) was found upon annealing. Certainly no metallographic evidence is currently available to support a precipitation-ageing hypothesis and the softening is therefore more consistent with the removal of embrittling elements from solution to the grain boundaries. There is some evidence which suggests that solid-solution hardening may occur in the case of beryllium alloys of higher iron content. Thus the differing degrees of dislocation interaction revealed in the electron micrographs (figs. 8 and 10) is indicative of different values for the stacking-fault energy of the 0.15 wt o/o alloy
126
o.
V.
T.
RANZETTA
in the annealed and homogenised conditions ; for example, quenching from 800" C resulted in higher values than were found after quenching this alloy from 920’ C. To some degree the differences are in accord with the mechanical properties of the two materials, since the latter condition was found by tensile testing to be more brittle in the temperature range 300” C to 600” C. This solid solution hardening may also be present in commercial-purity beryllium but to a lesser extent. Considering the grain boundaries - if these are brittle in homogenised metal, as both the mode of tensile fracture and the high tensile strength suggest, it is possible that they are strengthened on annealing. This could perhaps result from the removal of an embrittling thin film or ultra fine precipitates present at the boundary and the consequent formation of a coarser precipitate such as that observed in thin foils of annealed material. Should these precipitates act as dislocation sources, the ease with which dislocations sre generated could depend critically upon the precipitate distribution and the process may be more readily operative in annealed specimens. Consequently, in this material the propagation of plastic flow into neighbouring grains could be facilitated, with a corresponding relaxation of stress at the grain boundaries, and this would lead to the higher ductility and reduced tendency for intergranular failure which are observed experimentally. In conclusion, considerable experimental evidence is now being accumulated concerning the mechanical behaviour of beryllium and its related microstructure and the effect that thermal treatment may have upon this, but a comprehensive solution to the ductility problem in beryllium has not yet been achieved. For example, it is still uncertain whether or not very high purity beryllium would be ductile, although the present results on berylliumchromium alloys suggest that some improve-
AND
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ment of ductility in the high-temperature range ms,y be achieved by purification of the matrix. It would be particularly desirable to extend the present work to investigate the effeot that carbon, oxygen and nitrogen may have upon the properties of beryllium, and also to include some studies on ternary systems. Future work will be espeoislly with the properties of dislocations
concerned and their
distribution as a function both of temperature and of metal purity, and will be extended to lower temperatures, which are of particular interest to the metal technologist. Acknowledgements Acknowledgements sre due to the staff of the beryllium section under Dr. A. J. Martin, who provided the material and mechanical-test data and carried out the necessary heat treatments and specimen preparation. Thanks are also due to Professor J. Nutting, Dr. A. Moore and Mr. F. Morrow for helpful discussions. This work was carried out in the Research Metallurgy Branch, AWRE Aldermaston, under the direction of Mr. G. C. Ellis. References
1)
A. Moore, F. Morrow, V. D. Scott and D. A. Cheer, Conf. on Metallurgy London,
of Beryllium
(Inst. of Metals,
1961)
M. J. Whelan,
Proc. Roy
Sot.
A249 (1958) 114
Y 3)
D. McLean,
4,
C. Crussard, J. Plateau and G. Henry,
6)
Saclay, 1960) R. J. Teitel and
9
Univ.
Grain Boundaries
Press,
Y.
Adda,
Maurice
N.
(Oxford
1957) p. 116
of Grain Boundaries
Mining, Met.,
in Metals
(4th Metallurgy
M.
Petrol., Azam,
Cohen, Engrs. G.
Properties Colloquium,
Trans.
Am.
185 (1949)
Donze,
J.
Inst.
285
Mallen,
and M. Weisz, Compt. Rend.
Acad.
F. Sci.
‘1
(Paris), 254 (1962) 1052 A. J. Martin and G. C. Ellis, Conf. on Metallurgy
9
G. Donze,
9)
J. Mat. Nucl. 6 (1962) 137 A. Brown, F. Morrow and A. J. Martin, J. Less-
of Beryllium
Common
(Inst.
of Metals,
D. Dutilloy,
Metals
London,
R. le Hazif
3 (1961)
62
1961)
et Y.
Adda,