On the effects of vacuum annealing and carburizing on the ductility of coarse-grained molybdenum

Journal of the Less-Common Metals Elsevier Sequoia S.A., Lausanne - Printed in the Netherlands

245

ON THE EFFECTS OF VACUUM ANNEALING AND CARBURIZING THE DUCTILITY OF COARSE-GRAINED MOLYBDENUM

K. TSUYA

AND

ON

N. ARITOMI

Natioxal Research Institute for Metals, Tokyo (Japan) (Received

January

Igth, 1968)

SUMMARY

The cause of intergranular brittleness in molybdenum has been examined by studying the variation in ductility with vacuum annealing, carburizing, and decarburizing, using electron-beam-melted, coarse-grained specimens. The intergranular brittleness could not be reduced by vacuum annealing at zooo’C. A suitable carburizing treatment much improved the ductility; many carbides were observed on the intergranular surfaces in carburized ductile specimens, suggesting that the existence of carbides is not a main cause of the intergranular brittleness. Carburized ductile specimens were embrittled by decarburization in hydrogen. The improvement in ductility resulting from carburizing is thus attributed not only to the deoxidation during the carburizing process but also to the effects of carbon itself and to dispersion of carbides in matrices. From observations on ageing it is suggested that increase in carbon segregation at grain boundaries may contribute towards reducing intergranular brittleness.

IKTRODUCTIOX

Coarse-grained molybdenum is usually brittle and fractures intergranularly at a low stress at room temperature. The intergranular brittleness has been attributed by several investigators to the existence of intergranular oxides132 or carbides394, or to the segregation of interstitial impurities at grain boundariess. Also, it has been reported that the ductility of single crystal6 or wrought bar7 molybdenum is reduced by carbon. On the other hand, it has been reported8 that molybdenum ingots, which have been electron-beam melted and contain a suitable amount of carbon, are ductile and free from intergranular brittleness. The decrease in the oxygen content of the ingot with the amount of carbon added could not be determined, and it appeared that the carbon remaining in the ingot might have contributed to the strengthening of the grain boundaries. The present work was undertaken mainly to determine whether the improvement in ductility by the addition of carbon is attributable to the deoxidation effect of carbon, or to the direct effect of carbon itself. To this end the variation in ductility with vacuum annealing, carburizing, and decarburizing treatments has been studied using electron-beam-melted, coarse-grained specimens. J. Less-Common. Metals, 15 (1968)

K. TSL-YA, N. ARITOMI

246 SPECIMENS ASD EXPERIMENTAL METHODS

Specimens

Two types of specimen, ingot and wire, were used, Ingot specimens were made by cutting from an electron-beam-melted button ingot (about 3 cm diam. and I cm thick) so that one grain boundary was closely perpendicular to the tensile axis. The dimensions of the specimens were about 1.5 x 4 x 20 mm. Since the grain size of the ingot was so large, the specimens were actually bi-crystal specimens in tensile tests. For investigations other than the effect of grain size on tensile properties wire specimens having a bamboo-type structure* and measuring I mm x 50 mm were used. These were made from an electron-beam-melted ingot (about 16 mm x 200 mm) by swaging and drawing, followed by annealing at 2000°C for 30 min under vacuum. The purity of the molybdenum powder used as the raw material in the production of both ingots was better than 99.95% by weight. The button ingot contained less than IO p.p.m. oxygen, 20-30 p.p.m. carbon, and about 30 p.p.m. nitrogen. In the wire specimen, corresponding impurity contents were about IO p.p.m., IO p.p.m. and 30 p.p.m. Experimental

methods

After grinding and electropolishing, specimens were annealed or carburized. They were heated in a vacuum of 7 x 10-6-3 x IO- 5 mm Hg, indirectly by a tantalum heater or directly by passing through them an alternating current. In both cases the temperature was raised to 2ooo°C within 5 min. Temperatures, controlled within + 50°C about a fixed temperature, were measured by a W/W-Re thermocouple, an optical pyrometer and an automatic two-colour pyrometer. Figure I shows the cooling rate of a specimen in the tantalum heating furnace when the current was interrupted and

Fig. I. Cooling rate of a specimen in tantalum

heating furnace.

when the specimen was rapidly cooled by helium. Cooling in the direct-heating method was more rapid than depicted in Fig. I, but no difference in ductility due to the greater cooling rate could be detected. Carburization was carried out by coating a specimen with a suspension of graphite (spectrographic electrode of 99.999% purity) in benzine, followed by drying and heating under vacuum. Decarburization was accomplished by heating at rjoo”C in * This term refers to a structure in which all grain boundaries are closely perpendicular axis, as with the nodes of bamboo. An example of the structure is given by f\LLEh*. J. Less-Commo~t

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to the wire

I)I~CTII.IT\’ OF (:Or\RSE-(;R.IISEI)

>fOLYHI)ES~>I

z-t7

flowing hydrogen of 99.99890 purity and containing 0.032 mg/l of water, o.o003’J/, oxygen and 0.00199/, nitrogen. After annealing, or carburizing, specimens were again electropolished. Carburized specimens in particular were electropolished sufficiently to remove surface carbide layers. After examination of the structure under the optical microscope, tensile test.;, wcrc carried out in an Instron-type testing machine at an estension speed of 0.~7 mm ‘min at room temperature. The gauge length was about 8 mm for the ingot specimen and 30 mm for the wire specimen. The intergranular fracture surfaces were examined under the optical microscope and by electron microscopy.

Prelimivavy

esanli~iatio?~ usiq

ifgot specimen

The variation of grain boundary fracture stress of the ingot with vacuum annealing and carburizing treatments is summarized in Fig. 2. Most as-cast specimens fractured intergranularly and the grain boundary fracture stresses showed wide scatter,

I;ig. L. lcffects of vacuum anncaling and carburizing on the grain boundary fracture stress of ingot specimen. Blank circles show intergranular fracture, solid circles show transgranular fracture, and circles which are not completely solid show inter- and transgranular fracture. Triangles show the specimens which fractured after extension beyond the maximum load in the load-elongation curve. Figures give the per cent reduction in area and arrows indicate that the grain boundary fracture stress exceeds that at the point shown. Test temperature: Lj’(‘.

ranging from 2 to 33 kg/mm2. Specimens which had been annealed at 2ooo“C for I II and 8 h, respectively, followed by furnace cooling, fractured intergranularly and the grain boundary fracture stress did not differ greatly from that for As-cast specimens. Specimens which were vacuum annealed at 2000~ C for I h followed by helium quenching, showed a completely brittle intergranular fracture at a low stress. On the other J. Less-Common

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hand, the grain boundary fracture stress was clearly increased by carburizing at 2oooT for I h. The carburized specimens were very ductile and fractured by shear after showing a reduction in area of nearly 100%. Although almost all the grain boundaries in the as-cast specimens were very clean, precipitates were sometimes observed along some grain boundaries. Figure 3 shows an intergranular fracture surface containing a considerable amount of precipitates. This structure is considered to be eutectic, consisting of carbides, oxides, or nitrides with Lw-molybdenumin a grain boundary, where impurities are concentrated by the rapid solidification. After annealing at 2oooT for I h under vacuum, the precipitates disappeared completely. The intergranular surfaces of the annealed specimens were smooth and contained a considerable number of large voids, as shown in Fig. 4; there was no difference between the furnace-cooled and thehelium-quenched specimens.

Fig. 3. Intergranular pitates. (x 200)

fracture surface of an as-cast specimen containing

a large amount of preci-

Fig. 4. Intergranular fracture surface of a specimen annealed at zooo’C for I h under vacuum. Before annealing, the specimen contained precipitates as shown in Fig. 3. ( x zoo)

Fig. 5. Surface pattern of an as-carburized

specimen,

showing feathery carbides.

(X

zoo)

The solubilities of carbon, oxygen and nitrogen in molybdenum at 2000°C are about 0.018, 0.01 and 0.007 wt.%, respectivelys. These figures are much greater than the contents of these impurities in the specimens used. Hence, the change in intergranular pattern resulting from the vacuum annealing may be caused by the solution of the precipitates into the matrices and the suppression of re-precipitation by the fairly rapid cooling rate employed. The voids observed in the intergranular surfaces are considered to be gas holes which were formed by the gases produced on the decomposition of oxides or nitrides, or by the mutual reaction of oxides and carbides. The fractographic observations show that the grain boundary fracture stress does not corJ.

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relate with the inter-granular pattern, and suggest that the main factor causing intergranular brittleness is not the intergranular precipitates. This is in agreement with results previously reporteds. The surfaces of as-carburized specimens were covered with carbides. The surface carbides in Fig. 5 are feathery in appearance 10; this structure is usually observed on those parts of the specimens which were not coated with graphite powder. Where a large amount of graphite powder was used, dark-gray layers, possible of carbide having a large C/MO ratio, were observed. After these layers were eliminated by removing abaut 0.1 mm by electrapolishing, the structure containing MO& particles (Fig. 6 (a)), appeared. From the examination of sections, it was found that carbon had distributed uniformly through the specimen. Since the carburized specimens were too ductile to fracture intergranularly in the tensile test, intergranular fracture surfaces were obtained together with transgranular surfaces by an impact fracture. Figure 6 (b) shows the carbides on the intergranular surface of the carburized specimen.

Fig. 6. Structure of the ductile ingot specimen which was carburized (a) Eiectropolished microstructure. (b) Intergranular fracture surface.

at ZOOOT for

I

h. (X zoo)

These results indicate that grain boundaries are not markedly strengthened by vacuum annealing at 2000°C for 8 h, although the inter-granular pattern changes distinctly, while the carburizing treatment strengthens the grain boundaries and much improves the ductility. As, however, the ingot specimen is practically a bi-crystal, it is not suitable for a detailed investigation of the ductility change, since the shape of the stress-elongation curve varies greatly with the crystal orientation of the specimen. The bamboo-type structure specimens, containing 20-30 differently oriented crystals lying in the direction of the specimen axis, were used therefore for further experiments on vacuum annealing, carburizing, and decarburizing. It was expected that the ductility of these specimens might be greatly affected by the cohesive force between grain boundaries, and that the shape of the stress-elongation curve might not vary as much as that of the ingot specimens.

Figure 7 shows the effect of grain size on the tensile behaviour of wire specimens. The specimens, 0.3-0.4 mm in average grain diameter, which had been annealed at 15ooY for 30 min followed by furnace cooling, showed a yield point and fractured after developing elongation of about 30% with necking. These specimens were not of a bamboo-type structure. On the other hand, the specimens which had been annealed at J. Less-Common Metals, 15 (1968)

K. TSUYA,

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2oooT for 30 min followed by furnace cooling were of the bamboo-type structure, consisting of grains of about 1.1 mm in average diameter (average length of grain in the longitudinal direction). These specimens fractured intergranularly after an elongation of I-5:/,, without any yield point. Although WROSSKI et al.1’912 have reported that coarse-grained molybdenum is more ductile than fine grained, these results show

‘r ---Jo

rn’n

F.C

a

3

‘0

4

8III

12

’ 16 20 Elongation III,

24

28

32’0246

P/.)

Fig. 7. Typical stress-elongation curves for wire specimens annealed at 15oo~C and zooo3C, respectively. F.C. : furnace cooled, H.Q. : helium quenched. Test temperature: 25°C. Fig. 8. Ductile-to-brittle

transition behaviour

of bamboo-type-structure

specimen.

that the correlation between the ductility and the grain size is similar to that for other b.c.c. metals. The explanation of the disappearance of the yield point in the bambootype-structure specimen is considered to be as follows : the specimen contained 20-30 variously oriented crystals lying in the direction of the axis, and the moduli of elasticity of each crystal differ with each orientation 13. From analogy with the stressstrain curves for three differently oriented tungsten single crystalsl4, it is assumed that the molybdenum crystal favourably oriented to show a yield point, has the greatest modulus of elasticity. In ordinary polycrystal specimens, the deformation of each crystal is strongly restricted by adjoining crystals. However, since in the bamboo-type-structure specimen a large part of the surface of each crystal is not surrounded by other crystals, the independent deformation of each crystal is easier than in ordinary polycrystal specimens. Therefore, crystals favourably oriented to show a yield point may not be deformed plastically unless other crystals, not so favourably oriented, have been largely deformed plastically and hardened; this may offset the appearance of the yield point in the bamboo-type-structure specimen. Although a grain boundary acts essentially as a site of crack initiation and a fracture is not usually large grain size promotes intergranular cracking 15, intergranular observed in other b.c.c. metals having large solubilities for interstitial impurities, such as tantalum and niobium, in contrast with molybdenum. This suggests that the interfacial bonding force in molybdenum may be considerably lowered by some interstitial impurities. If so, it may be that grain growth reduces the grain boundary area per unit volume and increases the amount of impurity per unit area of the grain boundary, hence the fact that the grain bounpromoting intergranular brittleness 16. Furthermore, daries are closely perpendicular to the tensile axis may promote the intergranular brittleness of the bamboo-type-structure specimens. J.

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The ductile-to-brittle transition behaviour of the bamboo-type-structure specimen used as the starting material in the following experiments, is shown in Fig. 8, taking the reduction of area in the tensile test to be a measure of ductility. The intergranular fracture surfaces of the bamboo-type-structure specimens were smooth and very clean, but they sometimes contained fine precipitates considered to be carbides. The inter-granular patterns of specimens annealed at 2ooo’C for long periods or quenched in helium were similar in appearance.

On quenching in helium the specimens became very brittle and fractured intergranularly at a low stress within an apparently elastic range of the stress-elongation curve, as shown in Fig. 7. The embrittlement of the grain boundaries is not considered to be due to oxygen contamination of the helium because the grain boundaries are similarly embrittled by quenching in high-purity argon containing less than 0.1 p.p.m. oxygen, and because the intergranular patterns of the specimens quenched in helium are not different from those of furnace-cooled specimens. The residual thermal stresses in the quenched specimens, if any, seem to be too small to affect the ductility, and it is therefore suggested that the embrittlement may be caused by the decrease in the amount of carbon segregated at grain boundaries, as described below; this same phenomenon has been observed in iron’?. Variation

in ductility

on vacuum

anzlzealing

The grain sizes of specimens armealed at zooo°C for z h and 6 h, respectively, followed by furnace cooling, were of the same order as those of specimens annealed for 30 min. The specimens annealed for a long period also fractured intergranularly after a small amount of elongation, as shown in Fig. 7. This indicates that the grain boundaries in the bamboo-type-structure specimens are not strengthened by the solid state refining effect of the vacuum annealing. On the other hand, FEW AND MANNINCP have reported that the ductility of a sintered molybdenum wire, 0.04 in. diam. (about I mm), is improved by vacuum annealing at 3800°F (about 21oo"C), and that the improvement in ductility is marked when the temperature of the specimen is raised to the annealing temperature in a short time. In the present work the time necessary to reach the annealing temperature was under 5 min, and no significant variation in ductility caused by the difference in heating rate was found. Furthermore, the difference in annealing temperature between the two experiments does not result in different results. The conflict between these two experimental results may be sought in differences in the specimens used; i.e., those used by FEW AND IV~ANKING possibly contained a considerable amount of oxygen and were usually 0.025-o.035 mm in grain diameter, whereas those used in the present study were of a coarse-grained bamboo-type-structure and contained only a very small amount of oxygen.

Typical stress-elongation curves for specimens carburized under different conditions are summarized in Fig. 9. The ductility and fracture mode of the specimens carburized at 500°C for 30 min were not different from those of the starting material. After carburizing for 12 h, some of the specimens fractured by shear, showing a reduction in area of IOO%, while the others fractured intergranularly after showing some

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reduction in area. The intergranular fracture surface was similar to that of the starting material and it seems that carburization hardly occurred in the specimens. Carburizing at IOOO~Cfor 30 min does not improve the ductility. Specimens carburized for 2 h showed a somewhat larger elongation though the fracture was completely intergranular. Specimens carburized for 5 h showed a yield point and fractured by shear after a reduction in area of 100%. On the intergranular surface, as obtained by an impact fracture, globular precipitates 0.5-1 ,u in diameter, thought to be carbides, were observed under the electron microscope.

1000°C

Elongation

15OOT

2OOO~C

P/d

Fig. 9. Typical stress-elongationcurvesfor bamboo-type-structure specimens which were carburized by graphite coating followed by vacuum anealing at LOO”, IOOO*, rgoo" and zooo”C, respectively. Test temperature: 25°C.

Fig. IO. Intergranular fracture surfaces of ductile specimens which were carburized 5 min (a) and for 30 min (b), respectively. (x 200)

at 1500°C for

Specimens carburized at r500°C for 5, IO and 30 min, respectively, were very ductile and fractured by shear after a reduction in area of 100%. The yield stress and elongation per cent of these specimens increased with carburizing time. Intergranular carbides increase in amount with carburizing time, as shown in Fig. IO (a) and (b). Specimens carburized at zooo’C were brittle. The fracture consisted of interand transgranular surfaces on which large carbides were observed. The results show that the ductility of coarse-grained molybdenum can be greatly improved by a suitable carburizing treatment. However, from comparison with the results for the ingot specimen described above, it is considered that the limiting size and amount of carbide particles which are not harmful to the ductility may be largely dependent on the grain size and dimension of the specimen. J. Less-Common Metals, 15 (1968)

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To investigate whether the improvement in ductility ment was caused by the deoxidation during the carburizing of carbon itself, the variation in ductility on decarburizing (a) shows the typical stress-elongation curves for specimens zing at 15oo’C for IO min, and then decarburized at 15oo’C tively, in a hydrogen atmosphere. 40

(a)

due to carburizing treatprocess, or by the effect was examined. Figure II made ductile by carburifor 5, 8 and 15 h, respec-

(bf

Elongation

(%I

Fig. I I, Typical stress-elongation curves for (a) specimens which were decarburized in hydrogen at 1500°C after they had been made ductile by carburizing at 15oo’C for IO min., (b) specimens which were heated in hydrogen at 1500°C without carburizing before the treatment, and (c) specimens which were decarburized in hydrogen at r5oo”C for 8 and 15 h, respectively, followed by vacuum annealing at zooo°C for 30 min and then at x300°C for I h to precipitate carbides. Hours in the figure show the heating time in hydrogen. Test temperature: 25°C.

After decarburizing for 5 h, the yield point disappears and the tensile strength diminishes, but the specimens are still ductile and fracture by shear, showing a reduction in area of 100%. Specimens decarburized for 8 and 15 h, respectively, fractured intergranularly with reduced ductility but after some necking. The grain size is reduced on decarburizing, the average grain diameter falling from about 1.1 mm to around 0.7 mm. However, this reduction in grain diameter should contribute to the improvement in the ductility of the decarburized specimens. Specimens which had not been carburized were heated in the same hydrogen atmosphere at 1500°C for 5 and 8 h, respectively, and then tensile tested. The results are shown in Fig. I I (b). The treatment has tended to improve ductility ,although fracture is intergranular. The grain size in both instances is reduced by the annealing from about 1.1 mm to 0.8-0.9 mm average grain diameter; part of the improvement in ductility resulting from annealing in hydrogen may be associated with the decrease in grain size. Specimens which have not been carburized are not reduced in ductility by annealing in hydrogen, while the ductility of carburized specimens is reduced by the same treatment; this indicates that the ductility of carburized specimens is reduced by decarburization but not by contamination with impurities, particularly oxygen, contained in the hydrogen used. The improvement in ductility following carburizing may hence be attributed not only to the deoxidation during the carburizing process but also to the effect of carburization itself. After decarburizing, the intergranular fracture pattern of the specimen changes clearly to that shown in Fig. 12. The decrease in carbon content brought about by the decarburizing treatment was examined by observing the intergranular surface patJ. Less-Com.mon Metals,

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tern, after decarburizing, annealing under vacuum at zooo’C for 30 min to homogenize residual carbon, and finally annealing under vacuum at 13oo*Cfor x h to precipitate carbides. Figure 13 shows the intergranular surface of a specimen decarburized for 8 h and then annealed as described above. Angular and globular carbides are observed; the indication is that a considerable amount of carbon (estimated about O.OO~0.01 wt.?&) still remains in the specimen. The observations of the intergranular surfaces of the specimens decarburized for 15 h and then annealed as described above, indicate that the carbon content has been lowered to the level of that present before carburizing.

Fig. 12. Intergranular frakire surface of specimen which was decarburized hydrogen after carburizing at 15oo’C for IO min. (x zoo)

at rftoo”C

for 8 h in

l:ig. 13. Intergranular fracture surface of specimen which was decarburized in hydrogenat 15oo’C for 8 h, followed by vacuum annealing at zooo”C for 30 min and then at I 300% for I h to precipitate carbides. (x zoo)

Typical stress-elongation curves for the specimens which were decarburized and then heat treated to precipitate carbides, are shown in Fig. II (C). Comparing curves (C) and (A) it is apparent that the ductility of the specimens which have been decarburized for 8 h, and which contain a considerable amount of carbon, is not lowered by the treatment to precipitate carbides, but that the specimens decarburized for 1j h tend to be embrittled by the same treatment. As the grain size is not changed by the treatment, it is considered that the ductility of the specimens containing a small amount of carbon is lowered essentially by the treatment. Two possible mechanisms can be advanced to explain the improvement in ductility resulting from carburizing. Firstly, the cohesive force between grain boundaries is increased by carbon atoms segregated at those boundaries, as similarly applies in the case of iron”. The fact that the specimens decarburized for 15 h are embrittled by the treatment to precipitate carbides, although the specimens containing a considerable amount of carbon are not embrittled by the same treatment, suggests that the cohesive force between grain boundaries is reduced by the decrease in the amount of carbon dissolved around those boundaries. The observation that the ductility of specimens which are supersaturated with some interstitial impurities is improved by ageing, as described later, strongly supports this mechanism. On the other hand, carburized specimens containing a considerable amount of carbide show a yield point in the tensile test and are more ductile than aged specimens. This suggests for the second ductilizing mechanism that the dispersion of carbides in matrices may contribute to type yield point seen in the stressthe improvement in ductility; i.e., the rounded

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elongation curve probably indicates that carbides dispersed in matrices act as obstacles to the movement of dislocations, which may reduce the stress concentration at the grain boundaries and hence prevent intergranular fracture at low stresses. This mechanism may be supported by the observation that dispersion of ThOz in iron increases the cleavage fracture stress and improves the ductility at low temperaturels. These two mechanisms may be acting jointly to improve the ductility of carburized specimens. Imjwouement in ductility on ageing at 600°C It appears that the improvement in ductility of the specimens carburized at 500°C for 12 h (see Fig. 9) is not caused by carburization because the intergranular pattern is not changed by the carburizing treatment. To clarify this matter, tensile tests were carried out on specimens without graphite coating which had been annealed at 600°C for 12 h under vacuum. These specimens also fractured by shear, showing a reduction in area of roo”/o; the stress-elongation curves were similar to those for the specimens carburized at 500°C for 12 h. Hence, annealing at 500°C or 600°C for 12 h improves ductility whether or not carburization occurs, and it hardly does so at 500°C. It was also found that separately annealed specimens, the one at 600°C for 12 h in a reducing atmosphere of flowing CO and the other in an oxidizing atmosphere of flowing COZ, were both as ductile as those annealed under vacuum. The stress-elongation curves corresponding to all three treatments were similar in shape, suggesting that the improvement in ductility brought about by annealing is not attributable to the decrease in oxygen and carbon contents resulting from the mutual reaction of oxi des and carbides. Figure 14 depicts the solubilities of carbon, oxygen and nitrogen in molybdenum3. The levels of impurities contained in the bamboo-type-structure specimen used as starting material, are shown by the vertical lines, I, II and III.

1

*\

I \

24OOlm ‘\\ \( ,FL 2000..

: c

i

Few and ‘-Manning E-Perry&

01,

._a& and Rengstorfi

01 0

’

’

0.02 Carbonhvt. %)

Fig. 14. Tentative denum3.

0008 oxygen

0.016

wt.%)

Nitrogen

(wt.

%)

diagrams to illustrate s&id-solubility

of carbon, oxygen and nitrogen in molyb-

J. Less-Common Metals, 15 (1968)

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K. TSUYA,

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Since the specimens were cooled from 2ooo“C to room temperature at a fairly rapid rate (furnace cool in Fig. I), the lattice should be supersaturated with carbon, oxygen and nitrogen. When aged at 600°C for a long period, the specimens tended to approach an equilibrium condition. This may lead to the increase in the amount of segregation of these impurities at the grain boundaries. The activation energies for diffusion of carbon, oxygen and nitrogen in molybdenum are about 33,400,48,300 and 64,000 cal/mol respectivelyzo. Assuming that the diffusion coefficients (Do) of these atoms in molybdenum are all 0.01 cmz/sec20, the distances through which the carbon and oxygen move by diffusion during ageing at 600°C for 12 h are calculated to be about 0.02 mm and 0,0004mm, respectively; thus, carbon has the greatest tendency to segregate at grain boundaries. It is considered that the increase in the amount of carbon segregated at the grain boundaries may raise the cohesive force between the boundaries and so increase the ductility. It is believed that the mechanism of the improvement in ductility by the ageing treatment is not the same as that operating in the case of carburizing. The improvement in ductility by ageing may be explained in terms of the segregation of carbon at grain boundaries although the reason why such segregation improves the ductility is not clear. On the other hand, the dispersion of carbides may contribute to the improvement in ductility by carburizing. Such a mechanism, when supplemented by the increase in intergranular cohesive force by carbon, may be responsible for raising the ductility of carburized specimens beyond that of aged specimens. Although the aged specimens fracture by shear showing a reduction in area of 100% in the tensile test, they readily fracture intergranularly when repeatedly bent. On the contrary, specimens carburized at 15oo’C for IO min, infrequently fracture intergranularly by repeated bending and are more ductile than aged specimens. The thought that the improvement in ductility by ageing may be caused by the precipitation of fine carbides may be negated by the observation that no carbide was found on the intergranular surfaces of aged specimens, and such specimens do not show a yield point in the tensile test. These observations suggest that carbides are not present to a sufficient extent to act as barriers to the movement of dislocations within the metal. CONCLUSIONS

To investigate the cause of the intergranular brittleness in molybdenum, the variation in ductility with vacuum annealing, carburizing, and decarburizing was studied using electron-beam-melted, coarse-grained specimens. The results are as follows : (I) The grain boundaries in electron-beam-melted, coarse-grained molybdenum cannot be strengthened by the solid state refining effect of vacuum annealing at 2ooo°C. (2) A suitable carburizing treatment much improves the ductility and many carbides are observed on the intergranular surfaces in carburized, ductile specimens. This suggests that carbides are not a main cause of the intergranular brittleness in molybdenum. (3) Carburized, ductile specimens are embrittled bydecarburizationin hydrogen. The improvement in ductility on carburizing is attributed not only to the deoxidation J. Less-Common

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during the carburizing process but mainly to the effects of carbon itself and to the dispersion of carbides in matrices. (4) Specimens annealed at 2000°C under vacuum and cooled fairly rapidly to room temperature are brittle but become ductile on ageing at 5o0°-6o0°C for 12 h. This suggests that the increase in the amount of carbon segregated at grain boundaries may contribute towards reducing the intergranular brittleness. (5) It is considered that the cohesive force between grain boundaries in the electron-beam-nlelted, coarse-grained molybdenum increases with increasing carbon content, and that the effect of carbon is dependent on the amount segregated at grain boundaries but not on the total concentration of carbon. ACKNOWLEDGEMENT

The authors are indebted to Toho Kinzoku preparation of molybdenum wire.

Co. Ltd.,

Moji, Japan,

for the

REFERENCES

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