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Sinterability of fine molybdenum carbide powders obtained by a chemical vapour deposition method
Journal of the Less-Common Metals, 75 (1980) 11 - 21 0 Elsevier Sequoia S.A., Lausanne -Printed in the Netherlands
11
SINTERABILITY OF FINE MOLYBDENUM CARBIDE POWDERS OBTAINED BY A CHEMICAL VAPOUR DEPOSITION METHOD
J. HOJO, M. TAJIKA*
and A. KATO
Department of Applied Chemistry, Faculty of Engineering, 812 (Japan) (Received
Kyushu University, Fukuoka
October 30, 1979)
Summary The sinterabilities of fine MO& powders with a particle size of less than 0.3 pm produced by chemical vapour deposition (CVD) and commercial MosC powders with a particle size of about 2 I.trn were investigated at 900 1500 “C in hydrogen and at 1500 and 1750 “C in vacuum. Both types of powder gave sintered bodies with relative densities above 90% at 1500 “C. The powders produced by CVD showed a higher sinterability in vacuum than in hydrogen. It is suggested that free carbon retards the sintering of MO& particles. The higher sinterability of the powders chemically vapour deposited in vacuum is attributed to the removal of free carbon by the reaction of surface oxygen,
1. Introduction There are few studies on the sintering of submicron powders of transition metal carbides having extremely high melting points such as WC, MO& and Tic. The sintering of these carbides has generally utilized powders with particle sizes of greater than 2 I.tm and has required temperatures above 2000 “C [l, 21. We have previously reported the formation of ultrafine powders of molybdenum carbide (MO&) by a chemical vapour deposition (CVD) method [3 1. These powders are expected to show a high sinterability. In the present work we investigated their sintering behaviour and that of commercial MO& powders in hydrogen and in vacuum.
*Present address: Asahi Chemical Industry Co. Ltd., Fuji-shi 416, Japan.
12
2. Experimental 2.1. Materials Two kinds of MO& powders (designated A-MO& and B-MO&) with different particle sizes were produced by vapour phase reaction in the MoC14-CH4-Hs system at 1200 “C. The commercial MosC powders (designated C-MO&) were used as supplied (Mitsuwa’s Pure Chemicals). Electron micrographs of the MO& powders used are shown in Fig. 1 and their properties are summarized in Table 1. The particle size was determined by transmission electron microscopy (TEM). The specific surface area was measured by the Brunauer-Emmett-Teller method. The total carbon content and the nitrogen content* were determined by a combustion method. The free carbon content was determined by dissolving samples in an HF-HNOs mixture and analysing the carbon residue using a combustion method. The oxygen content was determined from the weight loss due to CO evolution during heating at 5 “C min-’ from 700 to 1100 “C in helium. The impurity metals detected by emission spectrochemical analysis are listed in Table 2. 2.2. Sin tering Powders, mixed with glycerine (<3 wt.%)** as plasticizer, were pressed at 4600 kgf cmT2 into pellets 10 mm in diameter and about 0.5 mm thick
A-Mo$
-
0,3 tlm
Fig. 1. TEM photographs
B-Flo2C -
0,3 um
C40~C
-
1 vrn
of the MO& powders.
*Nitrogen was used as the carrier gas for the MoC14 in the CVD method. **The glycerine in the compact was confirmed by thermogravimetric analysis to vaporize completely on heating at 10 “C min-’ to about 400 “C in hydrogen. The compact used in the wintering in vacuum was preheated to 500 “C in hydrogen to remove the glycerine.
13 TABLE 1 Properties of the MogC powders (wt.%)
Specific surface area (m2 g-l 1
Weight mean diameter (I-cm)
Chemical compositiona Total C
Free C
0
N
A-MogC
5.1
0.09
6.4
1.0
0.8
0.1
35 - 40
B-MO&
6.2
0.26
5.5
0.3
1.6
0.2
30 - 36
C-MogC
2.0
1.7
5.95b
O.llb
0.4
0.0
50 - 60
Sample
Bulk density of compact (%I
aBased on dry powders. bData given by the manufacturer.
TABLE 2 Emission spectrochemical
analysis of impurities in the MogC powders
Sample
Al
Mg
Co
Mn
Si
Fe
Cu
A-Mo~C
+
t
+
f
-
-
tr
B-MogC
+
+
f
tr
-
-
tr
C-Mo~C
tr
tr
+
-
t
-
tr
Ti
Co
Cr
Ni
W
+
+
f
-
+
The symbols show the intensities of spectral lines on films with the sequence + > + > - > tr. The orders of impurity content are as follows: +, 100 - 1000 ppm (Ti, Co), 10 - 100 ppm (Al, W); +, 10 - 100 ppm (Si, Mg, Ca, Mn, Cr); -, 10 - 100 ppm (Si), 1 - 10 ppm (Mn, Fe, Ni); tr, 1 - 10 ppm (Al, Mg, Mn, Cu).
with a single-action steel die. The pellets were set in a graphite box. For sintering in hydrogen (I), a recrystallized alumina tube (inner diameter 22 mm) was used as a vessel and an Sic resistance furnace as heater. The sample was moved from the low temperature zone (approximately 800 “C) to the centre of the furnace that had been heated to a given temperature to start the sintering. The hydrogen was deoxygenated in an activated copper column and was dried in a silica gel column and a liquid nitrogen trap. The gas flow rate was 100 ml min- ‘. For sintering in vacuum (II), a high frequency induction furnace with a molybdenum susceptor was used to heat the samples. The silica glass vessel used was evacuated to less than lo- 3 mmHg with an oil diffusion pump. In both experiments I and II the heating rate was about 100 “C mine1 from 800 “C to a given temperature. 2.3. Analysis of the siqtered body The linear shrinkage of the pellet diameter was measured with a cathetometer. The relative density (the X-ray density of MO& is 9.06 g cmW3 [4] )
14
and the porosities of open and closed pores were determined by an immersion method usingN,N-dimethylformamide (di5 = 0.945) (DMF) as solvent* (error in relative density, &HZ). The mi~ros~ct~e was observed by scanning electron microscopy (SEM). The sintered phase was identified by X-ray diffraction using Co Ka radiation. The Vickers microhardness was measured (error, &lo0 kgf mm-*).
3. Results 3.1. Sin tering in hydrogen 3.1.2. ~enaity of the sin tered body The variations with sintering time of the linear shrinkage of Mo2C sir&red at 1100 - 1500 “C in hydrogen are shown in Fig. 2. In both chemically vapour-deposited powders (A-Mo2C, B-Mo2C) and commercial powders (C-MO&), the sintered bodies shrink enormously during heating to a given ~mperat~e. The shrinkages at 0 min increase markedly with a rise in sintering temperature. The shrinkages of the B-Mo2C samples in particular change little with time during sintering at 1300 and 1500 “C. These facts indicate that densification of the compact occurs rapidly in the early stages of sintering. The linear shrinkage, relative density and porosities of A-, B- and C-MO& sintered for 30 min in hydrogen are shown against sintering temperature in Figs. 3 - 5. In A-MO&’ (Fig. 3) the density increases linearly with increasing temperature, being accompanied by a decrease in the open pore volume. The closed pore volume of sintered A-MO& is about 10% and changes little with temperature. In B-MO& (Fig. 4) the density increases markedly up to 1300 “C but only slightly between 1300 and 1500 “C. The closed pore volume of B-MO& sintered at low temperatures is as high as about 25% but this value decreases rapidly for B-MO& sintered at 1300 “C. In C-Mo2C (Fig. 5) the increase in the density becomes marked above 1300 “C. Sintering reduces the closed pore volume of C-Mo2C. B-Mo2C gives a denser sintered body than A-Mo2C although the particle size of B-Mo2C is larger than that of A-Mo2C. The density of the sintered B-MozC exceeds 90% at 1500 “C. In spite of the much larger initial particle size of C-MO&, the density of C-MO& sintered at 1500 “C is as high as that of B-MO& sintered at the same temperature. 3.1.2. Micrtxtructure The microstructures of Mo2C sintered in hydrogen are shown in Fig. 6. For A-Mo2C the bodies sintered at 1100 and 1300 “C consist of agglomerates
*The open pore porosity was determined by weighing the pellet impregnated with DMF in vacuum and the closed pore porosity was determined by weighing this impregnated pellet in DMF. From these weights the relative density was calculated.
15
0
a
I
I
I
I
15
30
45
60
Time
”
900
(min)
1100
1300
Temp.
(“C)
1500
Fig. 2. The variation with sintering time of the linear shrinkage of various MoaC samples sintered in hydrogen at various temperatures: A, A-Mo2C, 1100 “C; 0, B-MO&, 1100 “C; 0, B-Mo2C, 1300 “C; n, B-Mo2C, 1500 “C; 0, C-Mo,C, 1100 “C; 0, C-MoaC, 1300 “C; 0, C-MO& 1500 “C. Fig. 3. The variation with sintering temperature of the linear shrinkage AL/Lo (A), the relative density d, (a), the open pore porosity P, (0)and the closed pore porosity P, (0) of A-MoaC sintered in hydrogen for 30 min.
900
1100
Temp.
1300
(“Cl
1500
900
1100
1300
Temp.
(“C)
1500
Fig. 4. The variation with sintering temperature of the linear shrinkage, the relative density and the open and closed pore porosities of B-MO& sintered in hydrogen for 30 min. The symbols are the same as for Fig. 3. Fig. B.The variation with sintering temperature of the linear shrinkage, the relative density and the open and closed pore porasities of C-MozC sintered in hydrogen for 30 min. The symbols are the same aa for Fig. 3.
16
Fig. 6. SEM photographs of fracture surfaces of MozC samples sintered for 30 min in hydrogen at variaus temperatures: Al, A-Mo2C, 1100 “C; Aa, A-Mo2C, 1300 “C; Aa, A-Mo,C, 1500 “C; B1, B-Mo2C, 1100 “C; BP, B-Mo2C, 1300 “C!; B3, B-Mo2C, 1500 “C; Cl, C-MozC, 1100 “C; C2, C-MozC, 1300 “C; Ca, C-MozC, 1500 “C. (Scale, 2 pm long.)
of fine powders; gram growth becomes observable at 1500 “C. After 30 min at 1500 “C the grain size of A-Mo,C is about 0.5 pm. Grain growth of BMo,C is observable at the lower temperature of 1300 “C, indicating higher sinterability; at 1500 “C B-Mo,C has grains of about 1 Mm. Pore growth is observed together with the gram growth during the sinking of B-Mo2C. In C-Mo,C grain growth also becomes observable above 1300 “C; the grains become unusually coarse (about 8 pm) at 1500 “C. Sintered C-Mo,C contains large pores.
17
3.2. Sin Wing in vacuum The linear shrinkage, the relative density and the porosities of MO& sintered at 1500 and 1750 “C in vacuum are compared with the data for MO& sintered at 1500 “C in hydrogen in Table 3. For A-MO& and B-MO& the bodies sintered at 1500 “C are denser in vacuum than in hydrogen; in particular the density of B-MO& reaches 99%. An increase in temperature to 1750 “C has little effect on the density of A-MO& but lowers the density of B-Mo#. In contrast, for C-Mo2C the densities of bodies sintered in vacuum at both 1500 and 1750 “C are similar to the density of bodies sir&red at 1500 “C in hydrogen. The microstructures of MO& sintered in vacuum are shown in Fig. 7. At 1500 “C all samples of A-, B- and C-MO& form similar microstructures in vacuum and in hydrogen. Although there is little difference in the microstructures of A-MO& and C-Mo# sintered at 1500 and 1750 “C, enormous grain growth to 2 - 3 I_tm is observed in B-MO&! at 1750 “C; this may be responsible for the decrease in density. 3.3. X-ray analysis of sintered MO& Foreign phases appearing during sintering were investigated by X-ray diffraction of the pellet surface; such phases were detected only above 1500 “C and are listed in Table 4. Although the original powders are singlephase CY-MO~C,P-Mo2C appears during the sir&ring of all samples at 1500 “C in hydrogen. In vacuum, graphite appears at 1500 “C and graphite and q-MoC at 1750 “C in A-Mo2C! whereas in B-MozC metallic molybdenum appears at both 1500 and 1750 “C.
TABLE
3
A comparison of sintering data for Mo2C powders sintered in vacuum for 30 min with data for Mo2C sintered in hydrogen for 30 min SlllTZple
In hydrogen A-Mo2C B-Mo2C C-Mo2C
In vacuum A-Mo2C B-Mo2C C-Mo2C
Temperature W)
AL/Lo (%I
d, @I
PO (%I
PC @I
1500 1500 1500
21 27 13
87 92 96
2 1 0
11 7 4
1500 1750 1500 1750 1500 1750
21 22 29 28 14 14
94 95
99 89 96 97
5 5 1 11 3 2
AL/Lo, linear shrinkage; d,, relative density; P,, PC, open pore and closed pore porosities respectively.
18
Fig. 7. SEM photographs of fracture surfaces of MozC samples sintered for 30 min in vacuum at various temperatures: A,, A-MosC, 1500 “C; AZ, A-MozC, 1750 “C; Bl, B-Mo2C, 1500 “C; B2, B-Mo2C, 1750 “C; C1, C-Mo2C, 1500 “C; C2, C-Mo2C, 1750 “C. (Scale, 2 pm long.)
3.4. The hardness of sintered MozC The microhardness of the surface of pellets sintered at 1500 “C was measured; the results are compared with the relative densities in Table 5. In contrast with the densities, the hardnesses of A-Mo2C and B-MO& sintered in hydrogen are higher than the hardnesses after sintering in vacuum. C-Mo,C sintered in hydrogen also has higher hardness than C-Mo2C sintered in vacuum, in spite of there being little difference between the densities in both atmospheres. These results together with Table 4 suggest that fi-Mo,C has a higher hardness than cr-Mo2C and that graphite and metallic molybdenum segregate on the grain boundary to weaken the adhesion of Mo2C grains. The highest hardness in the present work is 1830 kgf mm- 2 for chemically vapour-deposited powders (B-Mo,C) sintered to a density of 92%; this is close to the literature value of 1800 kgf mm- 2 [ 51.
19
TABLE 4 Foreign phases appearing during the sintering of Mo$
powders 1750 “C in uacuum
Sample
1500 “C in hydrogen
1500 “C in uacuum
A-MO&
fl-Mo2C
Graphite
Graphite, Q-MoC
B-MO&
p-Mo2C
MO
MO
C-Mo2C
fl-Mo2C
n.d.
n.d.
The raw powders were CU-MozC. n.d., not detected.
TABLE 5 The Vickers microhardness Sample
and the relative density of Mo2C sintered at 1500 “C
Sin tered in hydrogen
Sintered
in vacuum
d, (%I
HV (kgf mmm2)
d, (%)
HV (kgf mrnm2)-
A-Mo2C
87
1690
94
1520
B-Mo2C
92
1830
99
1330
C-Mo2C
96
1250
96
960
4. Discussion 4.1. Particlesize effect In the sintering in hydrogen the densification of the compact of fine powders obtained by CVD (A-Mo,C, B-Mo,C) is more rapid at low temperatures than that for the densificatiop of commercial powders with a larger particle size. At 1500 “C, however, the density of sir&red C-Mo,C is as high as the density of B-Mo2C. The closed pore volumes of sintered A-Mo,C and B-Mo,C are greater than the corresponding volume of sintered C-Mo2C. As seen in Fig. 6, A-Mo,C and B-Mo,C form agglomerates of fine particles at low temperatures. The closed pores in both samples may stem mainly from pores included in the agglomerates. The higher closed pore porosities may be the reason that densification of compacts of the chemically vapourdeposited powders is not as rapid at high temperatures as is anticipated from their fineness. The fact that the sinterability of B-Mo2C is higher than that of A-Mo,C suggests that factors other than particle size influence the sinterability. As shown in Table 1, A-Mo2C has a higher content of free carbon than B-Mo2C and this may retard the sintering of Mo2C particles. C-Mo2C contains a lower content of free carbon than either A-Mo2C or B-Mo2C; this may account for the high sinterability of C-Mo2C with its larger particle size. The chemical
20
composition of the original MoaC powders is discussed in relation to the effect of atmosphere in the next section. 4.2. The effect of atmosphere The sinterabilities of A-MozC and B-MO&! are higher in vacuum than in hydrogen whereas that of C-MozC is the same in both atmospheres. A-MO& and B-MO& are characterized by higher contents of free carbon and oxygen; this difference in chemical composition may be responsible for the difference in the effect of atmosphere. As shown in Table 4, the foreign phase appearing above 1500 “C depends on the atmosphere and the sample. The appearance of &MO&, in addition to the original ~-MO& phase, at 1500 “C in hydrogen is consistent with the transition temperature of 1420 “C between the Q!and p phases [6]. In contrast, @-MO& is not detected after sintering in vacuum, indicating that this phase is entirely transformed to the (Yphase during cooling after sintering. In the phase diagram of the MO-C system [6], both the (Yand p phases are nonstoichiometric with a deficiency of carbon atoms and the p phase has a wider composition range than the cr phase*. It is therefore possible that more free carbon dissolves in the /3phase above 1500 “C and that the graphite appearing in A-MO& sintered in vacuum is produced by reprecipitation of the dissolved carbon when 0 transforms to (Yduring cooling. The q-MoC in A-MO& sintered at 1750 “C in vacuum can be explained by the fact that the n phase is formed above 1655 “C at compositions with more than 6.05 wt.% C [6]. The metallic moly~enum appearing in B-MO& sintered in vacuum may be produced by decarburization by surface oxygen which lowers the total carbon content below the lower limit of the composition range of cu-Mo,C; B-MO& has a high oxygen content compared with its free carbon content. We confirmed by gas chromatography that CO is liberated from MO& powders on heating above 700 “C in helium; this supports the suggestion that decarburization by surface oxygen occurs during sintering in vacuum. In contrast, the absence of a detectable amount of metallic molybdenum in B-MO& sintered in hydrogen is attributed to reduction of the surface oxide layer by hydrogen and carburization of the resulting molybdenum to MO& by free carbon and methane generated by hydrogenation of the graphite box used as the sample container. This is also supported by results obtained with a thermobalance, which show that the reduction of MOO, powders to molybdenum proceeds in the temperature range 300 - 800 “C with heating at 5 “C min- ’ in hydrogen and that the carburization of molybdenum powders with methane starts at about 750 “C. Methane at a partial pressure of 10e3 atm is also detected by gas chromatography in the hydrogen sintering atmosphere at 1100 - 1500 “C. From these results the effect of atmosphere on the sinterability of MozC powders can be summarized as follows. Although free carbon retards *The com~jtion of the CYphase is 5.56 - 5.70 wt.% C at 1000 “C and that of the #3 phase is 5.41 - 6.05 wt.% C.
21
the sintering of MO& particles as described in Section 4.1, more free carbon on the particle surface is removed by reaction with the surface oxide layer in vacuum than in hydrogen. This may cause the sinterabilities of A-MO& and B-MO& with high contents of free carbon to be higher in vacuum than in hydrogen. The small effect of atmosphere on the sinterability of C-Mo,C can be attributed to low contents of both free carbon and oxygen.
References L. Ramqvist, Powder Metall., 9 (1966) 26. M. S. Koval’chenko and Yu. I. Rogovoi, Sou. Powder Metall. Met. Ceram., 9 (1970) 1013. (Translated fromF’oroshk. Metall., 10 (12) (1970) 72.) J. Hojo, M. Tajika and A. Kato, J. Less-Common Met., 66 (1979) 151. L. E. Toth, Transition metal carbides and nitrides. In J. L. Margrave (ed.), Refractory Materials, Vol. 7, Academic Press, New York, 1971, p. 6. E. K. Storms, The refractory carbides. In J. L. Margrave (ed.), Refractory Materials, Vol. 2, Academic Press, New York, 1967, p. 134. E. Rudy, St. Windisch, A. J. Stosick and J. R. Hoffman, Trans. Metall. Sot. AIME, 239 (1967) 1247.
11
SINTERABILITY OF FINE MOLYBDENUM CARBIDE POWDERS OBTAINED BY A CHEMICAL VAPOUR DEPOSITION METHOD
J. HOJO, M. TAJIKA*
and A. KATO
Department of Applied Chemistry, Faculty of Engineering, 812 (Japan) (Received
Kyushu University, Fukuoka
October 30, 1979)
Summary The sinterabilities of fine MO& powders with a particle size of less than 0.3 pm produced by chemical vapour deposition (CVD) and commercial MosC powders with a particle size of about 2 I.trn were investigated at 900 1500 “C in hydrogen and at 1500 and 1750 “C in vacuum. Both types of powder gave sintered bodies with relative densities above 90% at 1500 “C. The powders produced by CVD showed a higher sinterability in vacuum than in hydrogen. It is suggested that free carbon retards the sintering of MO& particles. The higher sinterability of the powders chemically vapour deposited in vacuum is attributed to the removal of free carbon by the reaction of surface oxygen,
1. Introduction There are few studies on the sintering of submicron powders of transition metal carbides having extremely high melting points such as WC, MO& and Tic. The sintering of these carbides has generally utilized powders with particle sizes of greater than 2 I.tm and has required temperatures above 2000 “C [l, 21. We have previously reported the formation of ultrafine powders of molybdenum carbide (MO&) by a chemical vapour deposition (CVD) method [3 1. These powders are expected to show a high sinterability. In the present work we investigated their sintering behaviour and that of commercial MO& powders in hydrogen and in vacuum.
*Present address: Asahi Chemical Industry Co. Ltd., Fuji-shi 416, Japan.
12
2. Experimental 2.1. Materials Two kinds of MO& powders (designated A-MO& and B-MO&) with different particle sizes were produced by vapour phase reaction in the MoC14-CH4-Hs system at 1200 “C. The commercial MosC powders (designated C-MO&) were used as supplied (Mitsuwa’s Pure Chemicals). Electron micrographs of the MO& powders used are shown in Fig. 1 and their properties are summarized in Table 1. The particle size was determined by transmission electron microscopy (TEM). The specific surface area was measured by the Brunauer-Emmett-Teller method. The total carbon content and the nitrogen content* were determined by a combustion method. The free carbon content was determined by dissolving samples in an HF-HNOs mixture and analysing the carbon residue using a combustion method. The oxygen content was determined from the weight loss due to CO evolution during heating at 5 “C min-’ from 700 to 1100 “C in helium. The impurity metals detected by emission spectrochemical analysis are listed in Table 2. 2.2. Sin tering Powders, mixed with glycerine (<3 wt.%)** as plasticizer, were pressed at 4600 kgf cmT2 into pellets 10 mm in diameter and about 0.5 mm thick
A-Mo$
-
0,3 tlm
Fig. 1. TEM photographs
B-Flo2C -
0,3 um
C40~C
-
1 vrn
of the MO& powders.
*Nitrogen was used as the carrier gas for the MoC14 in the CVD method. **The glycerine in the compact was confirmed by thermogravimetric analysis to vaporize completely on heating at 10 “C min-’ to about 400 “C in hydrogen. The compact used in the wintering in vacuum was preheated to 500 “C in hydrogen to remove the glycerine.
13 TABLE 1 Properties of the MogC powders (wt.%)
Specific surface area (m2 g-l 1
Weight mean diameter (I-cm)
Chemical compositiona Total C
Free C
0
N
A-MogC
5.1
0.09
6.4
1.0
0.8
0.1
35 - 40
B-MO&
6.2
0.26
5.5
0.3
1.6
0.2
30 - 36
C-MogC
2.0
1.7
5.95b
O.llb
0.4
0.0
50 - 60
Sample
Bulk density of compact (%I
aBased on dry powders. bData given by the manufacturer.
TABLE 2 Emission spectrochemical
analysis of impurities in the MogC powders
Sample
Al
Mg
Co
Mn
Si
Fe
Cu
A-Mo~C
+
t
+
f
-
-
tr
B-MogC
+
+
f
tr
-
-
tr
C-Mo~C
tr
tr
+
-
t
-
tr
Ti
Co
Cr
Ni
W
+
+
f
-
+
The symbols show the intensities of spectral lines on films with the sequence + > + > - > tr. The orders of impurity content are as follows: +, 100 - 1000 ppm (Ti, Co), 10 - 100 ppm (Al, W); +, 10 - 100 ppm (Si, Mg, Ca, Mn, Cr); -, 10 - 100 ppm (Si), 1 - 10 ppm (Mn, Fe, Ni); tr, 1 - 10 ppm (Al, Mg, Mn, Cu).
with a single-action steel die. The pellets were set in a graphite box. For sintering in hydrogen (I), a recrystallized alumina tube (inner diameter 22 mm) was used as a vessel and an Sic resistance furnace as heater. The sample was moved from the low temperature zone (approximately 800 “C) to the centre of the furnace that had been heated to a given temperature to start the sintering. The hydrogen was deoxygenated in an activated copper column and was dried in a silica gel column and a liquid nitrogen trap. The gas flow rate was 100 ml min- ‘. For sintering in vacuum (II), a high frequency induction furnace with a molybdenum susceptor was used to heat the samples. The silica glass vessel used was evacuated to less than lo- 3 mmHg with an oil diffusion pump. In both experiments I and II the heating rate was about 100 “C mine1 from 800 “C to a given temperature. 2.3. Analysis of the siqtered body The linear shrinkage of the pellet diameter was measured with a cathetometer. The relative density (the X-ray density of MO& is 9.06 g cmW3 [4] )
14
and the porosities of open and closed pores were determined by an immersion method usingN,N-dimethylformamide (di5 = 0.945) (DMF) as solvent* (error in relative density, &HZ). The mi~ros~ct~e was observed by scanning electron microscopy (SEM). The sintered phase was identified by X-ray diffraction using Co Ka radiation. The Vickers microhardness was measured (error, &lo0 kgf mm-*).
3. Results 3.1. Sin tering in hydrogen 3.1.2. ~enaity of the sin tered body The variations with sintering time of the linear shrinkage of Mo2C sir&red at 1100 - 1500 “C in hydrogen are shown in Fig. 2. In both chemically vapour-deposited powders (A-Mo2C, B-Mo2C) and commercial powders (C-MO&), the sintered bodies shrink enormously during heating to a given ~mperat~e. The shrinkages at 0 min increase markedly with a rise in sintering temperature. The shrinkages of the B-Mo2C samples in particular change little with time during sintering at 1300 and 1500 “C. These facts indicate that densification of the compact occurs rapidly in the early stages of sintering. The linear shrinkage, relative density and porosities of A-, B- and C-MO& sintered for 30 min in hydrogen are shown against sintering temperature in Figs. 3 - 5. In A-MO&’ (Fig. 3) the density increases linearly with increasing temperature, being accompanied by a decrease in the open pore volume. The closed pore volume of sintered A-MO& is about 10% and changes little with temperature. In B-MO& (Fig. 4) the density increases markedly up to 1300 “C but only slightly between 1300 and 1500 “C. The closed pore volume of B-MO& sintered at low temperatures is as high as about 25% but this value decreases rapidly for B-MO& sintered at 1300 “C. In C-Mo2C (Fig. 5) the increase in the density becomes marked above 1300 “C. Sintering reduces the closed pore volume of C-Mo2C. B-Mo2C gives a denser sintered body than A-Mo2C although the particle size of B-Mo2C is larger than that of A-Mo2C. The density of the sintered B-MozC exceeds 90% at 1500 “C. In spite of the much larger initial particle size of C-MO&, the density of C-MO& sintered at 1500 “C is as high as that of B-MO& sintered at the same temperature. 3.1.2. Micrtxtructure The microstructures of Mo2C sintered in hydrogen are shown in Fig. 6. For A-Mo2C the bodies sintered at 1100 and 1300 “C consist of agglomerates
*The open pore porosity was determined by weighing the pellet impregnated with DMF in vacuum and the closed pore porosity was determined by weighing this impregnated pellet in DMF. From these weights the relative density was calculated.
15
0
a
I
I
I
I
15
30
45
60
Time
”
900
(min)
1100
1300
Temp.
(“C)
1500
Fig. 2. The variation with sintering time of the linear shrinkage of various MoaC samples sintered in hydrogen at various temperatures: A, A-Mo2C, 1100 “C; 0, B-MO&, 1100 “C; 0, B-Mo2C, 1300 “C; n, B-Mo2C, 1500 “C; 0, C-Mo,C, 1100 “C; 0, C-MoaC, 1300 “C; 0, C-MO& 1500 “C. Fig. 3. The variation with sintering temperature of the linear shrinkage AL/Lo (A), the relative density d, (a), the open pore porosity P, (0)and the closed pore porosity P, (0) of A-MoaC sintered in hydrogen for 30 min.
900
1100
Temp.
1300
(“Cl
1500
900
1100
1300
Temp.
(“C)
1500
Fig. 4. The variation with sintering temperature of the linear shrinkage, the relative density and the open and closed pore porosities of B-MO& sintered in hydrogen for 30 min. The symbols are the same as for Fig. 3. Fig. B.The variation with sintering temperature of the linear shrinkage, the relative density and the open and closed pore porasities of C-MozC sintered in hydrogen for 30 min. The symbols are the same aa for Fig. 3.
16
Fig. 6. SEM photographs of fracture surfaces of MozC samples sintered for 30 min in hydrogen at variaus temperatures: Al, A-Mo2C, 1100 “C; Aa, A-Mo2C, 1300 “C; Aa, A-Mo,C, 1500 “C; B1, B-Mo2C, 1100 “C; BP, B-Mo2C, 1300 “C!; B3, B-Mo2C, 1500 “C; Cl, C-MozC, 1100 “C; C2, C-MozC, 1300 “C; Ca, C-MozC, 1500 “C. (Scale, 2 pm long.)
of fine powders; gram growth becomes observable at 1500 “C. After 30 min at 1500 “C the grain size of A-Mo,C is about 0.5 pm. Grain growth of BMo,C is observable at the lower temperature of 1300 “C, indicating higher sinterability; at 1500 “C B-Mo,C has grains of about 1 Mm. Pore growth is observed together with the gram growth during the sinking of B-Mo2C. In C-Mo,C grain growth also becomes observable above 1300 “C; the grains become unusually coarse (about 8 pm) at 1500 “C. Sintered C-Mo,C contains large pores.
17
3.2. Sin Wing in vacuum The linear shrinkage, the relative density and the porosities of MO& sintered at 1500 and 1750 “C in vacuum are compared with the data for MO& sintered at 1500 “C in hydrogen in Table 3. For A-MO& and B-MO& the bodies sintered at 1500 “C are denser in vacuum than in hydrogen; in particular the density of B-MO& reaches 99%. An increase in temperature to 1750 “C has little effect on the density of A-MO& but lowers the density of B-Mo#. In contrast, for C-Mo2C the densities of bodies sintered in vacuum at both 1500 and 1750 “C are similar to the density of bodies sir&red at 1500 “C in hydrogen. The microstructures of MO& sintered in vacuum are shown in Fig. 7. At 1500 “C all samples of A-, B- and C-MO& form similar microstructures in vacuum and in hydrogen. Although there is little difference in the microstructures of A-MO& and C-Mo# sintered at 1500 and 1750 “C, enormous grain growth to 2 - 3 I_tm is observed in B-MO&! at 1750 “C; this may be responsible for the decrease in density. 3.3. X-ray analysis of sintered MO& Foreign phases appearing during sintering were investigated by X-ray diffraction of the pellet surface; such phases were detected only above 1500 “C and are listed in Table 4. Although the original powders are singlephase CY-MO~C,P-Mo2C appears during the sir&ring of all samples at 1500 “C in hydrogen. In vacuum, graphite appears at 1500 “C and graphite and q-MoC at 1750 “C in A-Mo2C! whereas in B-MozC metallic molybdenum appears at both 1500 and 1750 “C.
TABLE
3
A comparison of sintering data for Mo2C powders sintered in vacuum for 30 min with data for Mo2C sintered in hydrogen for 30 min SlllTZple
In hydrogen A-Mo2C B-Mo2C C-Mo2C
In vacuum A-Mo2C B-Mo2C C-Mo2C
Temperature W)
AL/Lo (%I
d, @I
PO (%I
PC @I
1500 1500 1500
21 27 13
87 92 96
2 1 0
11 7 4
1500 1750 1500 1750 1500 1750
21 22 29 28 14 14
94 95
99 89 96 97
5 5 1 11 3 2
AL/Lo, linear shrinkage; d,, relative density; P,, PC, open pore and closed pore porosities respectively.
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Fig. 7. SEM photographs of fracture surfaces of MozC samples sintered for 30 min in vacuum at various temperatures: A,, A-MosC, 1500 “C; AZ, A-MozC, 1750 “C; Bl, B-Mo2C, 1500 “C; B2, B-Mo2C, 1750 “C; C1, C-Mo2C, 1500 “C; C2, C-Mo2C, 1750 “C. (Scale, 2 pm long.)
3.4. The hardness of sintered MozC The microhardness of the surface of pellets sintered at 1500 “C was measured; the results are compared with the relative densities in Table 5. In contrast with the densities, the hardnesses of A-Mo2C and B-MO& sintered in hydrogen are higher than the hardnesses after sintering in vacuum. C-Mo,C sintered in hydrogen also has higher hardness than C-Mo2C sintered in vacuum, in spite of there being little difference between the densities in both atmospheres. These results together with Table 4 suggest that fi-Mo,C has a higher hardness than cr-Mo2C and that graphite and metallic molybdenum segregate on the grain boundary to weaken the adhesion of Mo2C grains. The highest hardness in the present work is 1830 kgf mm- 2 for chemically vapour-deposited powders (B-Mo,C) sintered to a density of 92%; this is close to the literature value of 1800 kgf mm- 2 [ 51.
19
TABLE 4 Foreign phases appearing during the sintering of Mo$
powders 1750 “C in uacuum
Sample
1500 “C in hydrogen
1500 “C in uacuum
A-MO&
fl-Mo2C
Graphite
Graphite, Q-MoC
B-MO&
p-Mo2C
MO
MO
C-Mo2C
fl-Mo2C
n.d.
n.d.
The raw powders were CU-MozC. n.d., not detected.
TABLE 5 The Vickers microhardness Sample
and the relative density of Mo2C sintered at 1500 “C
Sin tered in hydrogen
Sintered
in vacuum
d, (%I
HV (kgf mmm2)
d, (%)
HV (kgf mrnm2)-
A-Mo2C
87
1690
94
1520
B-Mo2C
92
1830
99
1330
C-Mo2C
96
1250
96
960
4. Discussion 4.1. Particlesize effect In the sintering in hydrogen the densification of the compact of fine powders obtained by CVD (A-Mo,C, B-Mo,C) is more rapid at low temperatures than that for the densificatiop of commercial powders with a larger particle size. At 1500 “C, however, the density of sir&red C-Mo,C is as high as the density of B-Mo2C. The closed pore volumes of sintered A-Mo,C and B-Mo,C are greater than the corresponding volume of sintered C-Mo2C. As seen in Fig. 6, A-Mo,C and B-Mo,C form agglomerates of fine particles at low temperatures. The closed pores in both samples may stem mainly from pores included in the agglomerates. The higher closed pore porosities may be the reason that densification of compacts of the chemically vapourdeposited powders is not as rapid at high temperatures as is anticipated from their fineness. The fact that the sinterability of B-Mo2C is higher than that of A-Mo,C suggests that factors other than particle size influence the sinterability. As shown in Table 1, A-Mo2C has a higher content of free carbon than B-Mo2C and this may retard the sintering of Mo2C particles. C-Mo2C contains a lower content of free carbon than either A-Mo2C or B-Mo2C; this may account for the high sinterability of C-Mo2C with its larger particle size. The chemical
20
composition of the original MoaC powders is discussed in relation to the effect of atmosphere in the next section. 4.2. The effect of atmosphere The sinterabilities of A-MozC and B-MO&! are higher in vacuum than in hydrogen whereas that of C-MozC is the same in both atmospheres. A-MO& and B-MO& are characterized by higher contents of free carbon and oxygen; this difference in chemical composition may be responsible for the difference in the effect of atmosphere. As shown in Table 4, the foreign phase appearing above 1500 “C depends on the atmosphere and the sample. The appearance of &MO&, in addition to the original ~-MO& phase, at 1500 “C in hydrogen is consistent with the transition temperature of 1420 “C between the Q!and p phases [6]. In contrast, @-MO& is not detected after sintering in vacuum, indicating that this phase is entirely transformed to the (Yphase during cooling after sintering. In the phase diagram of the MO-C system [6], both the (Yand p phases are nonstoichiometric with a deficiency of carbon atoms and the p phase has a wider composition range than the cr phase*. It is therefore possible that more free carbon dissolves in the /3phase above 1500 “C and that the graphite appearing in A-MO& sintered in vacuum is produced by reprecipitation of the dissolved carbon when 0 transforms to (Yduring cooling. The q-MoC in A-MO& sintered at 1750 “C in vacuum can be explained by the fact that the n phase is formed above 1655 “C at compositions with more than 6.05 wt.% C [6]. The metallic moly~enum appearing in B-MO& sintered in vacuum may be produced by decarburization by surface oxygen which lowers the total carbon content below the lower limit of the composition range of cu-Mo,C; B-MO& has a high oxygen content compared with its free carbon content. We confirmed by gas chromatography that CO is liberated from MO& powders on heating above 700 “C in helium; this supports the suggestion that decarburization by surface oxygen occurs during sintering in vacuum. In contrast, the absence of a detectable amount of metallic molybdenum in B-MO& sintered in hydrogen is attributed to reduction of the surface oxide layer by hydrogen and carburization of the resulting molybdenum to MO& by free carbon and methane generated by hydrogenation of the graphite box used as the sample container. This is also supported by results obtained with a thermobalance, which show that the reduction of MOO, powders to molybdenum proceeds in the temperature range 300 - 800 “C with heating at 5 “C min- ’ in hydrogen and that the carburization of molybdenum powders with methane starts at about 750 “C. Methane at a partial pressure of 10e3 atm is also detected by gas chromatography in the hydrogen sintering atmosphere at 1100 - 1500 “C. From these results the effect of atmosphere on the sinterability of MozC powders can be summarized as follows. Although free carbon retards *The com~jtion of the CYphase is 5.56 - 5.70 wt.% C at 1000 “C and that of the #3 phase is 5.41 - 6.05 wt.% C.
21
the sintering of MO& particles as described in Section 4.1, more free carbon on the particle surface is removed by reaction with the surface oxide layer in vacuum than in hydrogen. This may cause the sinterabilities of A-MO& and B-MO& with high contents of free carbon to be higher in vacuum than in hydrogen. The small effect of atmosphere on the sinterability of C-Mo,C can be attributed to low contents of both free carbon and oxygen.
References L. Ramqvist, Powder Metall., 9 (1966) 26. M. S. Koval’chenko and Yu. I. Rogovoi, Sou. Powder Metall. Met. Ceram., 9 (1970) 1013. (Translated fromF’oroshk. Metall., 10 (12) (1970) 72.) J. Hojo, M. Tajika and A. Kato, J. Less-Common Met., 66 (1979) 151. L. E. Toth, Transition metal carbides and nitrides. In J. L. Margrave (ed.), Refractory Materials, Vol. 7, Academic Press, New York, 1971, p. 6. E. K. Storms, The refractory carbides. In J. L. Margrave (ed.), Refractory Materials, Vol. 2, Academic Press, New York, 1967, p. 134. E. Rudy, St. Windisch, A. J. Stosick and J. R. Hoffman, Trans. Metall. Sot. AIME, 239 (1967) 1247.