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Effect of aging after carbon doping on the ductility of molybdenum
Journal of the Less-Common Metals, 75 (1980) 31- 42 0 Elsevier Sequoia S.A., Lausanne - Printed in the Netherlands
EFFECT OF AGING AFTER CARBON DOPING ON THE DUCTILITY MOLYBDENUM
YUTAKA
HIRAOKA,
31
OF
MASATOSHI OKADA and RYOJI WATANABE
National Research Institute for Metals, l-2-1, Sengen, Sakura-Mura, Niihari-Gun, Zbaragi 300-3 1 (Japan) (Received August 29,1979;
in revised form November 19, 1979)
Summary The low temperature tensile properties of undoped specimens and speeimens doped with up to 55 wt.ppm C before and after aging at 600 or 1500 “C were examined. Carbon doping from 14 to 30 wt.ppm induced a remarkable improvement in ductility, as shown by a sharp increase in uc and a sharp decrease in T,.Aging at 600 “C induced a slight decrease in ductility for all doped specimens, whilst aging at 1500 “C induced no monotonic changes in ductility. The latter result was interpreted by considering thermodynamics and dec~bu~~g during aging.
1. Introduction In our previous investigations [ 1, 2 ] it has been demonstrated that a small amount of carbon, up to 25 wt.ppm, remarkably improves the low temperature ductility of recrystallized sintered-molybdenum sheets and this beneficial effect of carbon remains after electron beam welding. Furthermore, such a small amount of carbon changes the fracture mode from intergranular to cleavage. These results have been explained mainly by an increase in the crack prop~ation stress through grain boundaries and/or the lattice, caused by the increase in carbon concentration at grain boundaries, However, the content or the distribution of carbon which is most effective in enhancing the low temperature ductility is still not clear. The purpose of the present work is to investigate in detail the contribution of the carbon distribution to the low temperature ductility of molybdenum. The mechanical properties of molybdenum sheets, which had been carbon doped up to 55 wt.ppm, were examined before and after aging at 600 or 1500 “C by means of tensile tests at temperatures from -100 to 70 “C. These two aging temperatures were chosen because the solubility of carbon in molybdenum is very small at 600 “C and is reasonably large at 1500 “C. So it is probable that there will be precipi~tion of carbon at 600 “C and resolution of carbides at 1500 “C.
32
2. Experimental procedures Commercial scored-moly~enum sheets (1 mm thick) were annealed at 1500 “C for 1 h under a vacuum of about 1.3 X 10T4 Pa and then were furnace cooled. Carbon doping of these recrystallized materials was carried out as described in our previous report [l] . Aging treatments were performed under a vacuum of about 1.3 X 10e4 Pa. The homogenizing conditions and the carbon contents of undoped and doped materials before and after aging are listed in Table 1. The oxygen and nitrogen contents were 24 ?I 2 wt.ppm and less than 5 wt.ppm respectively for all materials. Tensile test specimens with a gauge size of 14 mm long and 4 mm wide were cut from the undoped and doped materials. The tensile test specimens were pulled to fracture at a crosshead speed of 1 mm min- ’ at temperatures from -100 to 70 “C. Scanning electron microscopy (SEM) was applied to determine the fracture mode of the specimen, which was broken in the tensile test. The distribution of carbon at grain boundaries was also examined by SEM. TABLE 1 Chemical analysis of carbon content Symbol
Treatment
Undoped Doped at 1200 “C for Doped at 1400 “C for Doped at 1500 “C for Doped twice at 1500 20 min Doped three times at for 20 min
(wt.ppm)
Carbon content Undoped or doped
24 h aging at 600 “C
1 h aging at 1500 “C
190 min 30 min 20 min “C for
7 14 18 30 42
7 13 17 33 44
5 8 14 25 39
1500 “C
55
-
-
3. Results 3.1. Effects of carbon doping up to 55 wt.ppm
The tensile data in undoped and doped specimens before aging are listed in Table 2. The specimens doped with 14 - 55 wt.ppm C, in comparison with the undoped specimens, showed large increases in fracture strength at low temperatures but only slight increases in yield strength. The doped specimens also showed a large increase in elongation below -50 “C, although the specimens doped with 55 wt.ppm C showed a smaller elongation than the other doped specimens. Additions of more than 25 wt.ppm C induced no further improvement in ductility in agreement with our previous results [ 1, 21.
TABLE 2 Tensile data of undoped and doped specimens before aging Specimen
Temperature (“C)
Yield strength W’a)
Fracture strength (MPa)
Elongation (XI
Undoped
70 20 -20 -50 -100
187 288 405 521 fa
450 654 511 520 607
53.8 43.5 4.3 0.8 0
Doped with 14 wt.ppm C
70 20 -20 -50 -70 -100
307 408 529 598 673 809
512 500 602 791 888 819
70.6 45.7 42.6 25.7 19.2 1.3
Doped with 18 wt.ppm C
70 20 -20 -50 -80 -100
303 377 479 577 719 P
464 562 605 731 796 835
51.7 55.4 36.5 10.9 6.4 0
Doped with 30 wt.ppm C
70 20 -20 -50 -75 -100
250 337 438 551 671 P
448 574 678 828 838 792
52.0 53.8 37.8 25.3 9.6 0
Doped with 42 wt.ppm C
70 20 -20 -50 -80 -100
258 338 451 574 694 696
474 531 684 807 811 726
61.0 45.4 37.3 18.7 7.3 1.3
Doped with 55 wt.ppm C
20 -20 -50 -67 -90
358 474 568 660 fa
496 711 705 749 770
39.8 26.3 7.1 4.8 0.7
‘f, fractured within elastic limit,
3.2. Effect of aging at 600 “C for 24 h Tensile data for the undoped and doped specimens after aging at 600 “C for 24 h are listed in Table 3. The yield and fracture strengths of the doped specimens were decreased slightly by aging whilst those of the undoped specimens were almost unchanged. The elongation was decreased slightly by aging for the doped specimens.
34 TABLE 3 Tensile data of undoped and doped specimens after aging at 600 “C for 24 h Temperature
Specimen
WI
Yield strength WV
Fracture strength WW
Elongation (%I
Undoped
70 20 -20 -50 -100
188 287 408 523 fs
513 591 673 524 661
57.6 51.3 19.3 1.3 0.2
Doped with 14 wt.ppm C
70 20 -20 -50 -80 -100
267 353 458 550 709 $
491 503 695 729 765 792
59.9 46.7 37.7 14.6 4.8 0
Doped with 18 wt.ppm C
70 20 -20 -50 -80 -100
230 342 452 541 672 fa
451 553 574 795 762 738
51.6 52.7 35.2 21.8 4.9 0
Doped with 30 wt.ppm C
70 20 -20 -50 -80 -100
251 305 429 528 651 P
438 505 690 677 711 753
54.6 51.3 45.2 6.9 3.1 0.3
Doped with 42 wtppm
70 20 -20 -50 -80 -100
225 318 378 530 641 fa
465 511 734 814 817 687
53.3 44.0 39.5 22.9 8.1 0
C
af, fractured within elastic limit.
3.3. Effect of aging at 1500 “C for 1 h Tensile data for the undoped and doped specimens after aging at 1500 “C for 1 h are listed in Table 4. The yield and fracture strengths and the elongation showed no monotonic changes on aging in contrast with aging at 600 “C. The yield strength was decreased appreciably for the specimens doped with 14 wt.ppm C and slightly for the specimens doped with 18 wt.ppm C, whilst the strength was almost unchanged for the undoped specimens and the specimens doped with more than 30 wt.ppm C. The fracture strength decreased appreciably for the undoped specimens and the specimens doped with 14 wt.ppm C, whilst this strength was almost unchanged for the specimens doped with more than 18 wt.ppm C. The elongation at low tem-
35 TABLE 4 Tensile data of undoped and doped specimens after aging at 1500 “C for 1 h Specimen
Temperature W)
Yield strength WW
Fracture strength WW
Elongation @I
Undoped
70 20 -20 -50 -100
184 287 402
486 513 420 377 362
56.3 14.6 1.5 0 0
70 20 -20 -50 -80 -100
201 297 434 499 F
463 576 735 698 648 567
59.5 58.6 34.2 8.7 0.4 0
Doped with 18 wt.ppm C
70 20 -20 -50 -80 -100
262 355 436 606 685 P
449 523 600 828 827 790
51.2 45.1 42.1 28.0 8.9 0
Doped with 30 wt.ppm C
70 20 -20 -50 -80 -100
289 376 488 586 712 P
475 533 660 777 809 789
57.6 52.3 46.2 16.5 7.5 0
Doped with 42 wt.ppm C
70 20 -20 -50 -80 -100
278 358 448 553 685 f”
453 535 555 792 811 775
52.8 45.9 37.5 21.7 8.5 0
Doped with 14 wt.ppm C
F
af, fractured within elastic limit.
perature was decreased by aging for the undoped specimens and the specimens doped with 14 wt.ppm C, whilst it was almost unchanged for the specimens doped with more than 18 wt.ppm C. 3.4. Fractography of the undoped and doped specimens Using SEM the intergranular fracture area as a percentage of the total fracture area ‘(AIF/ATF X 100) was calculated for the undoped specimens and the doped specimens before and after aging. Undoped specimens typically show intergranular fractures (about 90% intergranular), whilst the doped specimens all show mainly cleavage fractures (20 - 30% intergranular). Aging at 600 “C for 24 h induced a slight increase in AIF/ATF X 100 for the
36
doped specimens. Aging at 1500 “C for 1 h induced a comparative increase in this ratio for the undoped specimens and the specimens doped with 14 wt.ppm C whilst it induced no obvious changes for the specimens doped with more than 18 wt.ppm C. No carbides were observed at grain boundaries for the undoped specimens and the specimens doped with 14 wt.ppm C, whilst carbides were recognized in the specimens doped with more than 18 wt.ppm C and their size or density increased with increasing total carbon content. Figure 1 shows scanning electron micrographs from undoped and doped specimens. After aging at 600 or 1500 “C no visible metallographical changes were recognized in these specimens, except for the specimen doped with 18 wt.ppm C; in this specimen, carbides were hardly observed after aging at 1500 “C for 1 h. 3.5. Estimation of ac and Tc For the ductile-brittle transition behaviour, a model proposed by Wronski et al. [ 31 can be employed. They postulate that T> Tc
(Jd-Y)
T< Tc
ON> (Tp
(2)
T=Tc
uN(=:(ly) = (Ip = (SC
(3)
<
Fig. 1. Scanning electron micropaphs carbides.
UP
(1)
of undoped and doped specimens; arrows indicate
37
where uv is the yield strength and uN, up, uc and Z’c are the crack nucleation stress, the crack propagation stress, the critical stress and the critical temperature respectively. The parameters uc and T, can be estimated by using the tensile data. uc is the stress that propagates microcracks through grain boundaries and/or the lattice at and around T, [ 1,2] . Tc is considered to be the ductile-brittle transition temperature in the present discussion. Figure 2 shows the two parameters uc and l/T, as functions of carbon content for the undoped specimens and the specimens doped with up to 55 wt.ppm C. Two conclusions are drawn from this figure. One is that a small amount of carbon (14 wt.ppm C) induces a sharp increase in uc and a decrease in T,, and then these values are almost constant up to 55 wt.ppm C. The other conclusion is that the doped specimens with no precipitates (14 wt.ppm C) show uc values as high as those for the doped specimens with precipitates (more than 18 wt.ppm C). Figures 3 and 4 show the effect of aging at 600 and 1500 “C on uc and l/T,. uc was unchanged for the undoped specimens and slightly decreased for the doped specimens by aging at 600 “C for 24 h. However, on aging at 1500 “C for 1 h, uc decreased appreciably for the undoped specimens and the specimens doped with 14 wt.ppm C and decreased slightly for the specimens doped with 18 wt.ppm C, whilst it was almost unchanged for the specimens doped with more than 30 wt.ppm C. The changes in l/T, were consistent with those in uc.
J OOk------
10
20
CARBON
30
LO
CONTENT
50
680
(wt.ppm)
Fig. 2. UC and l/!fc vs. the total carbon content: 0, specimens containing precipitates; 0, specimens without precipitates.
300
24h AGING t AT 6OO’C
_
BEFORE _ AGING
lh AGING AT 15OO*C
24 h AGlNG AT 6&C
_
BEFORE AGING
_
1 hAGlNG AT 1500%
I
50
Fig. 3. The effect of aging at 600 and 1500 “C on UC. Symbols 0 - 4 represent undoped and doped specimens with progressively increasing carbon content between 14 and 42 wt.ppm C (see Table 1). Fig. 4. The effect of aging at 600 and 1500 “C on l/T, or Tc. Symbols 0 - 4 represent undoped and doped specimens with progressively increasing carbon content between 14 and 42 wt.ppm C (see Table 1).
4. Discussion 4.1. Improvement in ductility by a small amount of carbon The ratio AIF/ATF X 100 is shown as a function of uc and l/T, for the undoped and doped specimens before and after aging in Fig. 5. The uc values nearly all lie on a straight line, and by extrapolation the maximum critical stress uc* can be obtained. uc* is the stress that propagates microcracks through the lattice, i.e. the cleavage fracture stress of a specimen with a constant grain size (about 0.020 mm), and is independent of the amount or the distribution of carbon. This is supported by other workers [4, 51. The ductile cleavage fracture stress is given by [6] (4) where rdc is the value of the effective surface energy for fracture at the tip of the crack at the instant that it propagates, E is Young’s modulus and 1 and lo/2 are the grain diameter and the distance from the tip of an arrested crack to the gram boundaries respectively. By neglecting 1a/2, rdc could be calculated to be 43 J m-* at the temperature Tc* = -113 “C (160 K) for I = 0.020 mm and u; = uc* = 940 MPa. This value of 7dc agreed roughly with the value of 30 - 40 J m-* obtained by extrapolation from the results of Wronski and Johnson [6] for the variation of Tdc with grain size at 195 K.
39
Tc (“Cl
-\ \ 0
I
I
L
5
\ T;
6
I/ Tc ( 10-3K’) (a)
(b)
Fig. 5. Plots of (a) l/Tc and (b) UC US.AIFIATF x 100 for undoped before and after aging: 0, without precipitates, before aging; 0, with aging; 0, without precipitates, after aging at 600 “C for 24 h; n, with aging at 600 “C for 24 h; A, without precipitates, after aging at 1500 precipitates, after aging at 1500 “C for 1 h.
and doped specimens precipitates, before precipitates, after “C for 1 h; A, with
By considering the relation between uc and AI F/AT F X 100 and the fact that the cleavage fracture stress is constant, it can be deduced that the fracture mode of the specimen should be a function only of the strength of the grain boundary or the stress that propagates microcracks through grain boundaries. From Fig. 2 or Fig. 5 it can also be suggested that the element that improves the ductility of molybdenum is the carbon in solid solution, not the carbide at the grain boundary. This is consistent with one of the two possible mechanisms advanced by Tsuya and Aritomi [ 71; the cohesive force at the grain boundary is enhanced by the increase in the amount of carbon dissolved around these boundaries. ‘I’c* is the minimum obtainable critical temperature and is obtained from the linear plot of Fig. 5. It follows from eqn. (3) that, whilst uc is nearly equal to up at and around Tc [l, 21, l/T, depends on both up and uy. Hence changes in l/T, may not be consistent with those in uc. In our present results, however, the changes in uv produced by carbon doping or aging after doping were much smaller than those in up or uc, so that l/T, showed similar behaviour to uc. Lastly we should mention the possibility proposed by Kumar and Eyre [8] that the effect of carbon is indirect, i.e. that the main role of carbon in improving ductility of molybdenum is to reduce the degree of segregation of oxygen to the grain boundaries. The significant differences between their specimens with bamboo-type coarse grains and about 8 at.ppm 0 and ours with relatively fine recrystallized grains and about 150 at.ppm 0 make it fairly difficult to compare their results with ours. In particular the high level of oxygen in our materials might prevent the application of the model pro-
40
posed by Kumar and Eyre to our results, since 150 at.ppm is much greater than the solubility limit of oxygen at 1500 “C (51 at.ppm [ 91). Thus it is probable that the grain boundaries are initially fully saturated with oxygen and that the oxygen segregation to these boundaries may not be reduced by additional doping at 1500 “C or less. This is also supported by a comparison of the present results with those obtained previously [ 21; the low temperature ductility of specimens with a much smaller oxygen content (60 at.ppm) was almost the same as that of specimens with 150 at.ppm 0 for undoped (7 wt.ppm C) and doped (24 - 30 wt.ppm C) material and no effect was obtained by reducing the oxygen content. We are therefore forced to conclude that carbon in solid solution at grain boundaries has a direct effect in improving the ductility at least in the present case. 4.2. Effect of aging on the low temperature ductility of doped specimens Assuming that the activation energy for diffusion and the diffusion coefficients De of carbon in molybdenum are 33 400 cal mol-’ and 0.01 cm2 s-l at 600 “C [lo], and 41000 cal mol-’ and 0.029 cm2 s-l at 1500 “C [ 111, respectively, the distance through which the carbon moves by diffusion during aging is calculated to be about 0.020 mm in 24 h at 600 “C!and about 0.30 mm in 1 h at 1500 “C. The solubility of carbon in molybdenum is less than 1 wt.ppm at 600 “C and about 50 wt.ppm at 1500 “C [ll, 121, provided that an equilibrium condition has been attained at each temperature. Considering these two points, the following processes are supposed to occur during aging at each temperature. At 600 “C grain boundary segregation or precipitation of carbon results in a decrease in the carbon dissolved in the lattice and/or at grain boundaries. In contrast, at 1500 “C resolution of carbides results in an increase in the carbon dissolved in the lattice and/or at grain boundaries. Such changes in the carbon distribution are expected to induce changes in uv and u c, and as a result changes in T, . The amount of carbon in the lattice determines the yield strength uy [13], which should thus decrease on aging at 600 “C and increase on aging at 1500 “C. Our results are consistent with the former supposition, but not with the latter; the inconsistency is discussed later. In our investigation uc slightly decreased for the doped specimens on aging at 600 “C. This supports our suggestion on the role of carbon, but is inconsistent with the result of Tsuya and Aritomi [7] who have reported that aging at 600 “C for 12 h improves the ductility. The conflict between these two experimental results may arise from the difference in the primary state of carbon distribution in the specimens used. Since the specimens of Tsuya and Aritomi had been rapidly quenched from 2000 “C, it is probable that the lattice was supersaturated with carbon. Hence grain boundary segregation of carbon but hardly any precipitation of carbides occurred on aging at 600 “C for 12 h, thus giving an increase in ductility. Our specimens, in contrast, had been slowly cooled from 1500 “C or less, and so it is more probable that the grain boundary segregation and, to some extent, precipitation occurred before aging. Hence only precipitation
41
proceeded on aging at 600 “C for 24 h, giving a decrease in ductility. In practice, however, strict comparison of our results with those of Tsuya and Aritomi is difficult, since their specimens with fairly large grains were not carbon doped and had different contents of carbon, oxygen and nitrogen from ours. However, the inconsistent results on uc or uy after aging at 1500 “C cannot be explained only by the thermodynamic considerations mentioned earlier and should be interpreted by dec~b~zation during aging as shown in Table 1. For the doped specimens with no precipitates (14 wt.ppm C) only decarburization occurs in the lattice and at grain boundaries and it induces the decrease in uy and aC. For the specimens with precipitates (more than 18 wt.ppm C), in contrast, the resolution of carbides and the decarburization occur concurrently in the lattice and at grain boundaries. These two effects offset each other, and as a result almost no visible changes are recognized in ciy or uc by aging at 1500 “C. In conclusion the carbon in solid solution is the cause of the grain boundary strengthening and not the carbide.
5. Conclusions (1) A small amount of carbon (14 wt.ppm C) induced a remarkable improvement in ductility, as shown by a sharp increase in ec and a sharp decrease in T,. These values were almost constant up to 55 wt.ppm C. (2) The doped specimens with no precipitates (14 wt.ppm C) showed ductility as great as the specimens with precipitates (more than 18 wt.ppm C). This suggests that the cause of the grain boundary strengthening is the carbon in solid solution at these boundaries. (3) Aging at 600 “C for 24 h induced a slight decrease in ductility for the doped specimens. (4) Aging at 1500 “C for 1 h induced no monotonic changes in ductility. The ductility was decreased appreciably by aging of the specimens with no precipitates (less than 14 wt.ppm C), whilst it was almost unchanged for the specimens with precipitates (more than 18 wt.ppm C). These results are explained by considering both the~odyn~ics and dec~b~zation during aging. The carbon in solid solution is the cause of the strengthening of the grain boundary and not the carbide itself,
The authors are indebted to Mr. F. Morito and Dr. T. Noda for their valuable discussions and also to Dr. T. Suzuki for the chemical analyses of carbon, oxygen and nitrogen.
42
References 1 2 3 4 5 6 7 8 9 10 11 12 13
Y. Hiraoka, F. Morito, M. Okada and R. Watanabe, J. Nucl. Mater., 78 (1978) 192. Y. Hiraoka, M. Okada and R. Watanabe, J. Nucl. Mater., 83 (1979) 305. A. S. Wronski, A. C. Chilton and E. M. Capron, Acta Metall., 7 (1969) 751. C. Crussard, R. Borione, J. Plateau, Y. Morillon and F. Maratray, J. Iron Steel Inst. London, 183 (1956) 146. H. Matsui and H. Kimura, Mater. Sci. Eng., 29 (1977) 241. A. S. Wronski and A. A. Johnson, Philos. Mag., 7 (1962) 213. K. Tsuya and N. Aritomi, J. Less-Common Met., 15 (1968) 245. A. Kumar and B. L. Eyre, Harwell Rep. AERE - R9330, 1979; Proc. R. Sot. London, Ser. A, 370 (1980) 431. S. C. Scrivastava and L. L. Seigle, Metall. Trans., 5 (1974) 49. G. W. Brock, Trans. Metall. Sot. AZME, 221 (1961) 1055. P. S. Rudman, Trans. Metall. Sot. AZME, 239 (1967) 1949. E. Gehardt, E. Fromm and U. Roy, 2. Metallkd., 57 (1966) 732. A. H. Cottrell and B. A. Bilby, Proc. R. Sot. London, Ser. A, 62 (1949) 49.
EFFECT OF AGING AFTER CARBON DOPING ON THE DUCTILITY MOLYBDENUM
YUTAKA
HIRAOKA,
31
OF
MASATOSHI OKADA and RYOJI WATANABE
National Research Institute for Metals, l-2-1, Sengen, Sakura-Mura, Niihari-Gun, Zbaragi 300-3 1 (Japan) (Received August 29,1979;
in revised form November 19, 1979)
Summary The low temperature tensile properties of undoped specimens and speeimens doped with up to 55 wt.ppm C before and after aging at 600 or 1500 “C were examined. Carbon doping from 14 to 30 wt.ppm induced a remarkable improvement in ductility, as shown by a sharp increase in uc and a sharp decrease in T,.Aging at 600 “C induced a slight decrease in ductility for all doped specimens, whilst aging at 1500 “C induced no monotonic changes in ductility. The latter result was interpreted by considering thermodynamics and dec~bu~~g during aging.
1. Introduction In our previous investigations [ 1, 2 ] it has been demonstrated that a small amount of carbon, up to 25 wt.ppm, remarkably improves the low temperature ductility of recrystallized sintered-molybdenum sheets and this beneficial effect of carbon remains after electron beam welding. Furthermore, such a small amount of carbon changes the fracture mode from intergranular to cleavage. These results have been explained mainly by an increase in the crack prop~ation stress through grain boundaries and/or the lattice, caused by the increase in carbon concentration at grain boundaries, However, the content or the distribution of carbon which is most effective in enhancing the low temperature ductility is still not clear. The purpose of the present work is to investigate in detail the contribution of the carbon distribution to the low temperature ductility of molybdenum. The mechanical properties of molybdenum sheets, which had been carbon doped up to 55 wt.ppm, were examined before and after aging at 600 or 1500 “C by means of tensile tests at temperatures from -100 to 70 “C. These two aging temperatures were chosen because the solubility of carbon in molybdenum is very small at 600 “C and is reasonably large at 1500 “C. So it is probable that there will be precipi~tion of carbon at 600 “C and resolution of carbides at 1500 “C.
32
2. Experimental procedures Commercial scored-moly~enum sheets (1 mm thick) were annealed at 1500 “C for 1 h under a vacuum of about 1.3 X 10T4 Pa and then were furnace cooled. Carbon doping of these recrystallized materials was carried out as described in our previous report [l] . Aging treatments were performed under a vacuum of about 1.3 X 10e4 Pa. The homogenizing conditions and the carbon contents of undoped and doped materials before and after aging are listed in Table 1. The oxygen and nitrogen contents were 24 ?I 2 wt.ppm and less than 5 wt.ppm respectively for all materials. Tensile test specimens with a gauge size of 14 mm long and 4 mm wide were cut from the undoped and doped materials. The tensile test specimens were pulled to fracture at a crosshead speed of 1 mm min- ’ at temperatures from -100 to 70 “C. Scanning electron microscopy (SEM) was applied to determine the fracture mode of the specimen, which was broken in the tensile test. The distribution of carbon at grain boundaries was also examined by SEM. TABLE 1 Chemical analysis of carbon content Symbol
Treatment
Undoped Doped at 1200 “C for Doped at 1400 “C for Doped at 1500 “C for Doped twice at 1500 20 min Doped three times at for 20 min
(wt.ppm)
Carbon content Undoped or doped
24 h aging at 600 “C
1 h aging at 1500 “C
190 min 30 min 20 min “C for
7 14 18 30 42
7 13 17 33 44
5 8 14 25 39
1500 “C
55
-
-
3. Results 3.1. Effects of carbon doping up to 55 wt.ppm
The tensile data in undoped and doped specimens before aging are listed in Table 2. The specimens doped with 14 - 55 wt.ppm C, in comparison with the undoped specimens, showed large increases in fracture strength at low temperatures but only slight increases in yield strength. The doped specimens also showed a large increase in elongation below -50 “C, although the specimens doped with 55 wt.ppm C showed a smaller elongation than the other doped specimens. Additions of more than 25 wt.ppm C induced no further improvement in ductility in agreement with our previous results [ 1, 21.
TABLE 2 Tensile data of undoped and doped specimens before aging Specimen
Temperature (“C)
Yield strength W’a)
Fracture strength (MPa)
Elongation (XI
Undoped
70 20 -20 -50 -100
187 288 405 521 fa
450 654 511 520 607
53.8 43.5 4.3 0.8 0
Doped with 14 wt.ppm C
70 20 -20 -50 -70 -100
307 408 529 598 673 809
512 500 602 791 888 819
70.6 45.7 42.6 25.7 19.2 1.3
Doped with 18 wt.ppm C
70 20 -20 -50 -80 -100
303 377 479 577 719 P
464 562 605 731 796 835
51.7 55.4 36.5 10.9 6.4 0
Doped with 30 wt.ppm C
70 20 -20 -50 -75 -100
250 337 438 551 671 P
448 574 678 828 838 792
52.0 53.8 37.8 25.3 9.6 0
Doped with 42 wt.ppm C
70 20 -20 -50 -80 -100
258 338 451 574 694 696
474 531 684 807 811 726
61.0 45.4 37.3 18.7 7.3 1.3
Doped with 55 wt.ppm C
20 -20 -50 -67 -90
358 474 568 660 fa
496 711 705 749 770
39.8 26.3 7.1 4.8 0.7
‘f, fractured within elastic limit,
3.2. Effect of aging at 600 “C for 24 h Tensile data for the undoped and doped specimens after aging at 600 “C for 24 h are listed in Table 3. The yield and fracture strengths of the doped specimens were decreased slightly by aging whilst those of the undoped specimens were almost unchanged. The elongation was decreased slightly by aging for the doped specimens.
34 TABLE 3 Tensile data of undoped and doped specimens after aging at 600 “C for 24 h Temperature
Specimen
WI
Yield strength WV
Fracture strength WW
Elongation (%I
Undoped
70 20 -20 -50 -100
188 287 408 523 fs
513 591 673 524 661
57.6 51.3 19.3 1.3 0.2
Doped with 14 wt.ppm C
70 20 -20 -50 -80 -100
267 353 458 550 709 $
491 503 695 729 765 792
59.9 46.7 37.7 14.6 4.8 0
Doped with 18 wt.ppm C
70 20 -20 -50 -80 -100
230 342 452 541 672 fa
451 553 574 795 762 738
51.6 52.7 35.2 21.8 4.9 0
Doped with 30 wt.ppm C
70 20 -20 -50 -80 -100
251 305 429 528 651 P
438 505 690 677 711 753
54.6 51.3 45.2 6.9 3.1 0.3
Doped with 42 wtppm
70 20 -20 -50 -80 -100
225 318 378 530 641 fa
465 511 734 814 817 687
53.3 44.0 39.5 22.9 8.1 0
C
af, fractured within elastic limit.
3.3. Effect of aging at 1500 “C for 1 h Tensile data for the undoped and doped specimens after aging at 1500 “C for 1 h are listed in Table 4. The yield and fracture strengths and the elongation showed no monotonic changes on aging in contrast with aging at 600 “C. The yield strength was decreased appreciably for the specimens doped with 14 wt.ppm C and slightly for the specimens doped with 18 wt.ppm C, whilst the strength was almost unchanged for the undoped specimens and the specimens doped with more than 30 wt.ppm C. The fracture strength decreased appreciably for the undoped specimens and the specimens doped with 14 wt.ppm C, whilst this strength was almost unchanged for the specimens doped with more than 18 wt.ppm C. The elongation at low tem-
35 TABLE 4 Tensile data of undoped and doped specimens after aging at 1500 “C for 1 h Specimen
Temperature W)
Yield strength WW
Fracture strength WW
Elongation @I
Undoped
70 20 -20 -50 -100
184 287 402
486 513 420 377 362
56.3 14.6 1.5 0 0
70 20 -20 -50 -80 -100
201 297 434 499 F
463 576 735 698 648 567
59.5 58.6 34.2 8.7 0.4 0
Doped with 18 wt.ppm C
70 20 -20 -50 -80 -100
262 355 436 606 685 P
449 523 600 828 827 790
51.2 45.1 42.1 28.0 8.9 0
Doped with 30 wt.ppm C
70 20 -20 -50 -80 -100
289 376 488 586 712 P
475 533 660 777 809 789
57.6 52.3 46.2 16.5 7.5 0
Doped with 42 wt.ppm C
70 20 -20 -50 -80 -100
278 358 448 553 685 f”
453 535 555 792 811 775
52.8 45.9 37.5 21.7 8.5 0
Doped with 14 wt.ppm C
F
af, fractured within elastic limit.
perature was decreased by aging for the undoped specimens and the specimens doped with 14 wt.ppm C, whilst it was almost unchanged for the specimens doped with more than 18 wt.ppm C. 3.4. Fractography of the undoped and doped specimens Using SEM the intergranular fracture area as a percentage of the total fracture area ‘(AIF/ATF X 100) was calculated for the undoped specimens and the doped specimens before and after aging. Undoped specimens typically show intergranular fractures (about 90% intergranular), whilst the doped specimens all show mainly cleavage fractures (20 - 30% intergranular). Aging at 600 “C for 24 h induced a slight increase in AIF/ATF X 100 for the
36
doped specimens. Aging at 1500 “C for 1 h induced a comparative increase in this ratio for the undoped specimens and the specimens doped with 14 wt.ppm C whilst it induced no obvious changes for the specimens doped with more than 18 wt.ppm C. No carbides were observed at grain boundaries for the undoped specimens and the specimens doped with 14 wt.ppm C, whilst carbides were recognized in the specimens doped with more than 18 wt.ppm C and their size or density increased with increasing total carbon content. Figure 1 shows scanning electron micrographs from undoped and doped specimens. After aging at 600 or 1500 “C no visible metallographical changes were recognized in these specimens, except for the specimen doped with 18 wt.ppm C; in this specimen, carbides were hardly observed after aging at 1500 “C for 1 h. 3.5. Estimation of ac and Tc For the ductile-brittle transition behaviour, a model proposed by Wronski et al. [ 31 can be employed. They postulate that T> Tc
(Jd-Y)
T< Tc
ON> (Tp
(2)
T=Tc
uN(=:(ly) = (Ip = (SC
(3)
<
Fig. 1. Scanning electron micropaphs carbides.
UP
(1)
of undoped and doped specimens; arrows indicate
37
where uv is the yield strength and uN, up, uc and Z’c are the crack nucleation stress, the crack propagation stress, the critical stress and the critical temperature respectively. The parameters uc and T, can be estimated by using the tensile data. uc is the stress that propagates microcracks through grain boundaries and/or the lattice at and around T, [ 1,2] . Tc is considered to be the ductile-brittle transition temperature in the present discussion. Figure 2 shows the two parameters uc and l/T, as functions of carbon content for the undoped specimens and the specimens doped with up to 55 wt.ppm C. Two conclusions are drawn from this figure. One is that a small amount of carbon (14 wt.ppm C) induces a sharp increase in uc and a decrease in T,, and then these values are almost constant up to 55 wt.ppm C. The other conclusion is that the doped specimens with no precipitates (14 wt.ppm C) show uc values as high as those for the doped specimens with precipitates (more than 18 wt.ppm C). Figures 3 and 4 show the effect of aging at 600 and 1500 “C on uc and l/T,. uc was unchanged for the undoped specimens and slightly decreased for the doped specimens by aging at 600 “C for 24 h. However, on aging at 1500 “C for 1 h, uc decreased appreciably for the undoped specimens and the specimens doped with 14 wt.ppm C and decreased slightly for the specimens doped with 18 wt.ppm C, whilst it was almost unchanged for the specimens doped with more than 30 wt.ppm C. The changes in l/T, were consistent with those in uc.
J OOk------
10
20
CARBON
30
LO
CONTENT
50
680
(wt.ppm)
Fig. 2. UC and l/!fc vs. the total carbon content: 0, specimens containing precipitates; 0, specimens without precipitates.
300
24h AGING t AT 6OO’C
_
BEFORE _ AGING
lh AGING AT 15OO*C
24 h AGlNG AT 6&C
_
BEFORE AGING
_
1 hAGlNG AT 1500%
I
50
Fig. 3. The effect of aging at 600 and 1500 “C on UC. Symbols 0 - 4 represent undoped and doped specimens with progressively increasing carbon content between 14 and 42 wt.ppm C (see Table 1). Fig. 4. The effect of aging at 600 and 1500 “C on l/T, or Tc. Symbols 0 - 4 represent undoped and doped specimens with progressively increasing carbon content between 14 and 42 wt.ppm C (see Table 1).
4. Discussion 4.1. Improvement in ductility by a small amount of carbon The ratio AIF/ATF X 100 is shown as a function of uc and l/T, for the undoped and doped specimens before and after aging in Fig. 5. The uc values nearly all lie on a straight line, and by extrapolation the maximum critical stress uc* can be obtained. uc* is the stress that propagates microcracks through the lattice, i.e. the cleavage fracture stress of a specimen with a constant grain size (about 0.020 mm), and is independent of the amount or the distribution of carbon. This is supported by other workers [4, 51. The ductile cleavage fracture stress is given by [6] (4) where rdc is the value of the effective surface energy for fracture at the tip of the crack at the instant that it propagates, E is Young’s modulus and 1 and lo/2 are the grain diameter and the distance from the tip of an arrested crack to the gram boundaries respectively. By neglecting 1a/2, rdc could be calculated to be 43 J m-* at the temperature Tc* = -113 “C (160 K) for I = 0.020 mm and u; = uc* = 940 MPa. This value of 7dc agreed roughly with the value of 30 - 40 J m-* obtained by extrapolation from the results of Wronski and Johnson [6] for the variation of Tdc with grain size at 195 K.
39
Tc (“Cl
-\ \ 0
I
I
L
5
\ T;
6
I/ Tc ( 10-3K’) (a)
(b)
Fig. 5. Plots of (a) l/Tc and (b) UC US.AIFIATF x 100 for undoped before and after aging: 0, without precipitates, before aging; 0, with aging; 0, without precipitates, after aging at 600 “C for 24 h; n, with aging at 600 “C for 24 h; A, without precipitates, after aging at 1500 precipitates, after aging at 1500 “C for 1 h.
and doped specimens precipitates, before precipitates, after “C for 1 h; A, with
By considering the relation between uc and AI F/AT F X 100 and the fact that the cleavage fracture stress is constant, it can be deduced that the fracture mode of the specimen should be a function only of the strength of the grain boundary or the stress that propagates microcracks through grain boundaries. From Fig. 2 or Fig. 5 it can also be suggested that the element that improves the ductility of molybdenum is the carbon in solid solution, not the carbide at the grain boundary. This is consistent with one of the two possible mechanisms advanced by Tsuya and Aritomi [ 71; the cohesive force at the grain boundary is enhanced by the increase in the amount of carbon dissolved around these boundaries. ‘I’c* is the minimum obtainable critical temperature and is obtained from the linear plot of Fig. 5. It follows from eqn. (3) that, whilst uc is nearly equal to up at and around Tc [l, 21, l/T, depends on both up and uy. Hence changes in l/T, may not be consistent with those in uc. In our present results, however, the changes in uv produced by carbon doping or aging after doping were much smaller than those in up or uc, so that l/T, showed similar behaviour to uc. Lastly we should mention the possibility proposed by Kumar and Eyre [8] that the effect of carbon is indirect, i.e. that the main role of carbon in improving ductility of molybdenum is to reduce the degree of segregation of oxygen to the grain boundaries. The significant differences between their specimens with bamboo-type coarse grains and about 8 at.ppm 0 and ours with relatively fine recrystallized grains and about 150 at.ppm 0 make it fairly difficult to compare their results with ours. In particular the high level of oxygen in our materials might prevent the application of the model pro-
40
posed by Kumar and Eyre to our results, since 150 at.ppm is much greater than the solubility limit of oxygen at 1500 “C (51 at.ppm [ 91). Thus it is probable that the grain boundaries are initially fully saturated with oxygen and that the oxygen segregation to these boundaries may not be reduced by additional doping at 1500 “C or less. This is also supported by a comparison of the present results with those obtained previously [ 21; the low temperature ductility of specimens with a much smaller oxygen content (60 at.ppm) was almost the same as that of specimens with 150 at.ppm 0 for undoped (7 wt.ppm C) and doped (24 - 30 wt.ppm C) material and no effect was obtained by reducing the oxygen content. We are therefore forced to conclude that carbon in solid solution at grain boundaries has a direct effect in improving the ductility at least in the present case. 4.2. Effect of aging on the low temperature ductility of doped specimens Assuming that the activation energy for diffusion and the diffusion coefficients De of carbon in molybdenum are 33 400 cal mol-’ and 0.01 cm2 s-l at 600 “C [lo], and 41000 cal mol-’ and 0.029 cm2 s-l at 1500 “C [ 111, respectively, the distance through which the carbon moves by diffusion during aging is calculated to be about 0.020 mm in 24 h at 600 “C!and about 0.30 mm in 1 h at 1500 “C. The solubility of carbon in molybdenum is less than 1 wt.ppm at 600 “C and about 50 wt.ppm at 1500 “C [ll, 121, provided that an equilibrium condition has been attained at each temperature. Considering these two points, the following processes are supposed to occur during aging at each temperature. At 600 “C grain boundary segregation or precipitation of carbon results in a decrease in the carbon dissolved in the lattice and/or at grain boundaries. In contrast, at 1500 “C resolution of carbides results in an increase in the carbon dissolved in the lattice and/or at grain boundaries. Such changes in the carbon distribution are expected to induce changes in uv and u c, and as a result changes in T, . The amount of carbon in the lattice determines the yield strength uy [13], which should thus decrease on aging at 600 “C and increase on aging at 1500 “C. Our results are consistent with the former supposition, but not with the latter; the inconsistency is discussed later. In our investigation uc slightly decreased for the doped specimens on aging at 600 “C. This supports our suggestion on the role of carbon, but is inconsistent with the result of Tsuya and Aritomi [7] who have reported that aging at 600 “C for 12 h improves the ductility. The conflict between these two experimental results may arise from the difference in the primary state of carbon distribution in the specimens used. Since the specimens of Tsuya and Aritomi had been rapidly quenched from 2000 “C, it is probable that the lattice was supersaturated with carbon. Hence grain boundary segregation of carbon but hardly any precipitation of carbides occurred on aging at 600 “C for 12 h, thus giving an increase in ductility. Our specimens, in contrast, had been slowly cooled from 1500 “C or less, and so it is more probable that the grain boundary segregation and, to some extent, precipitation occurred before aging. Hence only precipitation
41
proceeded on aging at 600 “C for 24 h, giving a decrease in ductility. In practice, however, strict comparison of our results with those of Tsuya and Aritomi is difficult, since their specimens with fairly large grains were not carbon doped and had different contents of carbon, oxygen and nitrogen from ours. However, the inconsistent results on uc or uy after aging at 1500 “C cannot be explained only by the thermodynamic considerations mentioned earlier and should be interpreted by dec~b~zation during aging as shown in Table 1. For the doped specimens with no precipitates (14 wt.ppm C) only decarburization occurs in the lattice and at grain boundaries and it induces the decrease in uy and aC. For the specimens with precipitates (more than 18 wt.ppm C), in contrast, the resolution of carbides and the decarburization occur concurrently in the lattice and at grain boundaries. These two effects offset each other, and as a result almost no visible changes are recognized in ciy or uc by aging at 1500 “C. In conclusion the carbon in solid solution is the cause of the grain boundary strengthening and not the carbide.
5. Conclusions (1) A small amount of carbon (14 wt.ppm C) induced a remarkable improvement in ductility, as shown by a sharp increase in ec and a sharp decrease in T,. These values were almost constant up to 55 wt.ppm C. (2) The doped specimens with no precipitates (14 wt.ppm C) showed ductility as great as the specimens with precipitates (more than 18 wt.ppm C). This suggests that the cause of the grain boundary strengthening is the carbon in solid solution at these boundaries. (3) Aging at 600 “C for 24 h induced a slight decrease in ductility for the doped specimens. (4) Aging at 1500 “C for 1 h induced no monotonic changes in ductility. The ductility was decreased appreciably by aging of the specimens with no precipitates (less than 14 wt.ppm C), whilst it was almost unchanged for the specimens with precipitates (more than 18 wt.ppm C). These results are explained by considering both the~odyn~ics and dec~b~zation during aging. The carbon in solid solution is the cause of the strengthening of the grain boundary and not the carbide itself,
The authors are indebted to Mr. F. Morito and Dr. T. Noda for their valuable discussions and also to Dr. T. Suzuki for the chemical analyses of carbon, oxygen and nitrogen.
42
References 1 2 3 4 5 6 7 8 9 10 11 12 13
Y. Hiraoka, F. Morito, M. Okada and R. Watanabe, J. Nucl. Mater., 78 (1978) 192. Y. Hiraoka, M. Okada and R. Watanabe, J. Nucl. Mater., 83 (1979) 305. A. S. Wronski, A. C. Chilton and E. M. Capron, Acta Metall., 7 (1969) 751. C. Crussard, R. Borione, J. Plateau, Y. Morillon and F. Maratray, J. Iron Steel Inst. London, 183 (1956) 146. H. Matsui and H. Kimura, Mater. Sci. Eng., 29 (1977) 241. A. S. Wronski and A. A. Johnson, Philos. Mag., 7 (1962) 213. K. Tsuya and N. Aritomi, J. Less-Common Met., 15 (1968) 245. A. Kumar and B. L. Eyre, Harwell Rep. AERE - R9330, 1979; Proc. R. Sot. London, Ser. A, 370 (1980) 431. S. C. Scrivastava and L. L. Seigle, Metall. Trans., 5 (1974) 49. G. W. Brock, Trans. Metall. Sot. AZME, 221 (1961) 1055. P. S. Rudman, Trans. Metall. Sot. AZME, 239 (1967) 1949. E. Gehardt, E. Fromm and U. Roy, 2. Metallkd., 57 (1966) 732. A. H. Cottrell and B. A. Bilby, Proc. R. Sot. London, Ser. A, 62 (1949) 49.