High thermal conductive silicon nitride sintered body and method of producing the same

A high thermal conductive silicon nitride sintered body contains: 2.0-7.5% by weight of a rare earth element in terms of the amount of an oxide thereof; at most 0.3% by weight of Li, Na, K, Fe, Ca, Mg, Sr, Ba, Mn and B as impurity cationic elements in terms of total amount thereof; and, if necessary, at most 2.0% by weight of alumina and/or at most 2.0% by weight of aluminum nitride, and comprises a beta-phase type silicon nitride crystal and a grain boundary phase. The silicon nitride sintered body has a thermal conductivity of at least 60 W/m.multidot.K. Optionally, the sintered body further contains 0.2-3.0% by weight of at least one compound selected from the group consisting of the oxides, carbides, nitrides, silicides and borides of Ti, Zr, Hf, V, Nb, Ta, Cr, Mo and W. The sintered body has a porosity of at most 1.5% by volume, a thermal conductivity of at least 60 W/m.multidot.K, and a three-point bending strength of at least 80 kg/mm.sup.2 at a room temperature. The sintered body achieves high thermal conductivity and good heat-radiating characteristics, as well as the high-strength characteristics generally inherent in silicon nitride sintered bodies.

BACKGROUND OF THE INVENTION 
1. Field of the invention 
The present invention relates to a high thermal conductive silicon nitride 
sintered body and a method of producing the same. More particularly, the 
invention relates to a high thermal conductive silicon nitride sintered 
body which achieves high thermal conductivity and good heat-radiating 
characteristics, as well as the high strength characteristics generally 
inherent in silicon nitride, so as to be suitable for semiconductor 
substrate, various types of radiator plates, etc., and a method of 
producing the high thermal conductive silicon nitride sintered body. 
2. Description of the Related Art 
Ceramic sintered bodies containing silicon nitride as a main component have 
strong heat resistance. They resist temperatures as high as 1000.degree. 
C. or higher. Silicon nitride ceramic sintered bodies also have strong 
thermal shock resistance due to their low thermal expansivity. Because of 
these characteristics, silicon nitride ceramic sintered bodies are 
expected to be widely used as high-temperature structural materials, most 
of which are currently made of heat-resistant super alloys. In fact, 
silicon nitride ceramic sintered bodies are already used for high-strength 
heat-resistant components and parts of, for example, gas turbines, engines 
or steel making machines. Further, because of their high corrosion 
resistance to metal, some silicon nitride ceramic sintered bodies are 
applied to melt-resistant material for molten metal. Still further, 
because of their high abrasion resistance, some silicon nitride ceramic 
sintered bodies are applied to or tested for cutting tools or sliding 
parts such as bearings. 
Various sintering compositions for silicon nitride ceramic sintered bodies 
are known, for example: silicon nitride-yttrium oxide-aluminum oxide 
system; silicon nitride-yttrium oxide-aluminum oxide-aluminium nitride 
system; and silicon nitride-yttrium oxide-aluminum oxide-oxide of 
titanium, magnesium or zirconium. 
The oxides of rare earth elements, such as yttrium oxide (Y.sub.2 O.sub.3) 
in the sintering compositions listed above, have been widely used as 
sintering assistant agents. Such rare earth element oxides enhance the 
sintering characteristics of sintering materials and, therefore, achieve 
high density and high strength of the sintered products (sintered bodies). 
According to the conventional art, silicon nitride sintered bodies are 
generally mass-produced as follows. After a sintering assistant agent as 
mentioned above is added to the powder of silicon nitride, the mixture is 
molded to form a compact. Then, the compact is sintered in a sintering 
furnace at about 1600.degree.-1850.degree. C. for a predetermined period 
of time followed by cooling in the furnace. 
However, though the silicon nitride sintered body produced by the 
conventional method achieves high mechanical strengths such as toughness, 
the thermal conductivities thereof are significantly lower than those of 
aluminum nitride (AlN) sintered bodies, beryllium oxide (BeO) sintered 
bodies or silicon carbide (SIC) sintered bodies. Therefore, conventional 
silicon nitride sintered bodies are unsuitable for electronic materials, 
such as semiconductor substrates, that need good heat-radiating 
characteristics. Accordingly, the use of silicon nitride sintered body is 
thus limited. 
Aluminum nitride sintered bodies have high thermal conductivity and low 
thermal expansivity, compared with other ceramic sintered bodies. Aluminum 
nitride sintered bodies are widely used as packaging materials or 
materials of circuit base boards for semiconductor chips, which have been 
progressively improved in operational speed, output power, variety of 
functions and size. However, no conventional aluminum nitride sintered 
bodies achieve sufficiently high mechanical strengths. 
Therefore, there is a growing need for a ceramic sintered body having both 
high thermal conductivity and high strength. 
SUMMARY OF THE INVENTION 
Accordingly, an object of the present invention is to provide a silicon 
nitride sintered body having a high thermal conductivity and, therefore, 
good heat-radiating characteristics, as well as the high strength 
characteristics generally inherent in silicon nitride sintered body, and a 
method of producing the silicon nitride sintered body. 
To achieve the above object, the present inventor studied the effects of 
the types of silicon nitride powder, sintering assistant agent and 
additives, the amounts thereof used, and the sintering conditions on the 
characteristics of the final products, that is, the sintered bodies, by 
performing experiments. 
The experiments provided the following findings. 
A silicon nitride sintered body having both high strength and high thermal 
conductivity can be obtained by: adding certain amounts of a rare earth 
element and aluminum ingredients, such as aluminum nitride and alumina, to 
a highly-pure fine powder of silicon nitride; molding to form a compact 
and degreasing the compact; maintaining the compact at a predetermined 
high temperature for a certain period of time to sinter the compact so as 
to enhance the density thereof; and then gradually cooling the sintered 
body at a certain rate. In short, the above method significantly enhances 
the heat conductivity of a silicon nitride sintered body. 
Further, formation of a glass phase (amorphous phase) in the grain boundary 
phase is effectively suppressed by using a highly pure powder of silicon 
nitride containing significantly reduced amounts of oxygen and impurity 
cationic elements, and preparing a silicon nitride molded compact having a 
reduced thickness before sintering. Thereby, a silicon nitride sintered 
body having a high thermal conductivity of 60 W/m.multidot.K or higher can 
be obtained even if only a rare earth element, but no aluminum 
ingredients, is added to a silicon nitride material powder. 
Still further, a silicon nitride sintered body having a significantly 
enhanced thermal conductivity as well as high strength can be obtained by: 
adding certain amounts of a rare earth element and at least one compound 
selected from the group consisting of the oxides, carbides, nitrides, 
silicides and borides of Ti, Zr, Hf, V, Nb, Ta, Cr, Mo and W, and, if 
necessary, alumina and/or aluminum nitride, to a highly-pure fine powder 
material of silicon nitride; molding to form a compact and degreasing the 
compact; maintaining the compact at a predetermined high temperature for a 
certain period of time to sinter the compact so as to enhance the density 
thereof; and then gradually cooling the sintered body at a certain rate. 
Further, the grain boundary phase in the structure of a silicon nitride 
sintered body is changed from an amorphous phase to a phase including 
crystal phases by gradually cooling the sintered body at a rate of 
100.degree. C. per hour or lower while controlling the cooling rate, 
thereby achieving both high strength and high thermal conductivity. If a 
sintered body in a sintering furnace is cooled simply by switching off the 
furnace as performed according to the conventional method, the cooling 
rate is rather high, that is, about 400.degree.-800.degree. C. per hour. 
The present invention has been achieved on the basis of the above findings. 
One aspect of the present invention provides a high thermal conductive 
silicon nitride sintered body containing: 2.0-7.5% by weight of a rare 
earth element in terms of the amount of an oxide thereof; and at most 0.3% 
by weight of Li, Na, K, Fe, Ca, Mg, Sr, Ba, Mn and B as impurity cationic 
elements, and having a thermal conductivity of at least 60 W/m.multidot.K. 
Another aspect of the present invention provides a high thermal conductive 
silicon nitride sintered body containing: 2.0-7.5% by weight of a rare 
earth element in terms of the amount of an oxide thereof; at most 2.0% by 
weight of aluminum compound calculated as alumina; at most 0.3% by weight 
of Li, Na, K, Fe, Ca, Mg, Sr, Ba, Mn and B as impurity cationic elements 
in terms of total amount thereof, and comprising a silicon nitride crystal 
and a grain boundary phase. 
The silicon nitride sintered body of the invention may contain at most 2.0% 
by weight of aluminum nitride, instead of the alumina. Further, the 
silicon nitride sintered body may contain at most 2.0% by weight of 
aluminum in terms of the amount of alumina, or at most 2.0% by weight of 
aluminum nitride there being at most 2% by weight of aluminum nitride and 
aluminum compound calculated as alumina as a tool amount there of. 
A further aspect of the present invention provides a high thermal 
conductive silicon nitride sintered body containing: 2.0-7.5% by weight of 
a rare earth element in terms of the amount of an oxide thereof; 0.2-3.0% 
by weight of at least one compound selected from the group consisting of 
the oxides, carbides, nitrides, silicides and borides of Ti, Zr, Hf, V, 
Nb, Ta, Cr, Mo and W; at most 0.3% by weight of Li, Na, K, Fe, Ca, Mg, Sr, 
Ba, Mn and B as impurity cationic elements in terms of total amount 
thereof; and, if necessary, at most 2.0% by weight of aluminum in terms of 
the amount of alumina and/or at most 2.0% by weight of aluminum nitride, 
and comprising a silicon nitride crystal and a grain boundary phase. 
It is preferred that the ratio of the area of a crystal compound phase 
formed in the grain boundary phase to the area of the grain boundary phase 
be at least 20%. 
A still further aspect of the present invention provides a method of 
producing a high thermal conductive silicon nitride sintered body, 
comprising the steps of: forming a compact by molding a mixture obtained 
by adding 2.0-7.5% by weight of a rare earth element in terms of the 
amount of an oxide thereof, and at most 2.0% by weight of alumina and/or 
at most 2.0% by weight of aluminum nitride, to a silicon nitride powder 
which contains at most 1.7% by weight of oxygen, at most 0.3% by weight of 
Li, Na, K, Fe, Ca, Mg, Sr, Ba, Mn and B as impurity cationic elements in 
terms of total amount thereof, and at least 90% by weight of alpha-phase 
type silicon nitride, and which has an average grain size of at most 0.8 
.mu.m; degreasing the compact; sintering the compact at a temperature of 
1800.degree.-2000.degree. C. while pressurizing the atmosphere around the 
compact to form a sintered body; and cooling the sintered body at a 
cooling rate of at most 100.degree. C. per hour until the temperature is 
reduced to a point at which a liquid phase formed of the rare earth 
element during the sintering step solidifies. 
A further aspect of the present invention provides a method of producing a 
high thermal conductive silicon nitride sintered body, comprising the 
steps of: forming a compact by molding a mixture obtained by adding 
2.0-7.5% by weight of a rare earth element in terms of the amount of an 
oxide thereof, 0.2-3.0% by weight of at least one compound selected from 
the group consisting of the oxides, carbides, nitrides, silicides and 
borides of Ti, Zr, Hf, V, Nb, Ta, Cr, Mo and W, and, if necessary, 
0.1-2.0% by weight of alumina and/or 0.1-2.0% by weight of aluminum 
nitride, to a silicon nitride powder which contains at most 1.7% by weight 
of oxygen, at most 0.3% by weight of Li, Na, K, Fe, Ca, Mg, Sr, Ba, Mn and 
B as impurity cationic elements in terms of total amount thereof, and at 
least 90% by weight of alpha-phase type silicon nitride, and which has an 
average grain size of at most 0.8 .mu.m; degreasing the compact; sintering 
the compact at a temperature of 1800.degree.-2000.degree. C. while 
pressurizing the atmosphere around the compact; and cooling the sinter at 
a cooling rate of at most 100.degree. C. per hour until the temperature is 
reduced to a point at which a liquid phase formed of the rare earth 
element during the sintering step solidifies. 
The method of the present invention achieves a high thermal conductive 
silicon nitride sintered body having both good mechanical characteristics 
and good thermal conductive characteristics, more specifically, a porosity 
of at most 1.5% by volume, a thermal conductivity of at least 60 
W/m.multidot.K, and a three-point bending strength of at least 80 
kg/mm.sup.2 at a room temperature. 
To achieve good sintering characteristics, high strength and high thermal 
conductivity of the product, the silicon nitride powder which is used in 
the method of the invention and contained as a main component in the 
sintered body of the invention contains at most 1.7%, preferably, 
0.5-1.5%, by weight of oxygen at most 0.3%, preferably, 0.2% or less, by 
weight of Li, Na, K, Fe, Ca, Mg, Sr, Ba, Mn and B as impurity cationic 
elements in terms of total amount thereof, and at least 90%, more 
preferably, 93% or more, by weight of alpha-phase type silicon nitride 
having good sintering characteristics, and, further the powder has fine 
grains, that is, an average grain size of at most 0.8 .mu.m, more 
preferably, about 0.4-0.6 .mu.m. 
The use of a fine powder of silicon nitride having an average grain size of 
0.8 .mu.m or less facilitates forming a dense sintered body having a 
porosity of 1.5% or less without requiring a large amount of a sintering 
assistant agent, and therefore reduces the potential adverse effect of a 
sintering assistant agent on the thermal conductivity of the sintered 
body. 
The total amount of the impurity cationic elements contained in the 
sintered body of the invention, that is, Li, Na, K, Fe, Ca, Mg, Sr, Ba, Mn 
and B, is limited to at most 0.3% by weight in order to ensure the 
sintered body a thermal conductivity of 60 W/m.multidot.K, because the 
impurity cationic elements adversely affect the thermal conductivity of 
the sintered body. The use of a silicon nitride powder containing at least 
90% by weight of alpha-phase type silicon nitride, which has better 
sintering characteristics than a beta-phase type, facilitates producing a 
high-density sintered body. 
Examples of the rare earth element to be contained as a sintering assistant 
agent in a silicon nitride powder are Y, La, Sc, Pr, Ce, Nd, Dy, Ho and 
Gd. Such a rare earth element may be contained in a silicon nitride powder 
in the form of an oxide thereof or a substance which is changed into an 
oxide thereof during the sintering process. One or more kinds of such 
oxide or substance may be contained in a silicon nitride powder. Among 
them, yttrium oxide (Y.sub.2 O.sub.3) is particularly preferred. Such a 
sintering assistant agent reacts with the silicon nitride powder so as to 
form a liquid phase and thereby serves as a sintering promoter. 
The amount of a sintering assistant agent to be contained in the silicon 
nitride powder (material powder) needs to be within a range of 2.0-7.5% by 
weight in terms of the amount of an oxide thereof. If the amount is less 
than 2.0% by weight, the sintered body fails to achieve a sufficiently 
high density and, therefore, the strength and thermal conductivity thereof 
are reduced to undesired levels. If the amount is more than 7.5% by 
weight, an excessively large portion of the grain boundary phase is 
formed, thereby reducing the thermal conductivity and strength of the 
sintered body. The more preferred range of the amount of a sintering 
assistant agent is 3.0-6.0% by weight. 
Further, according to the present invent ion, alumina (Al.sub.2 O.sub.3) 
assists the effect of a sintering promoter, that is, a rare earth element. 
Alumina provide a particularly great assisting effect if a pressurizing 
sintering process is employed. The amount of alumina to be contained in a 
silicon nitride powder is limited to at most 2.0% by weight because if the 
amount is greater than 2.0% by weight, an excessively large portion of the 
grain boundary phase is formed, or alumina starts dissolving into the 
silicon nitride and, therefore, reduces the thermal conductivity of the 
sintered body. Further, the amount of alumina contained is preferably 
within a range of 0.1-2.0% by weight because if it is less than 0.1% by 
weight, the sintered body fails to achieve a sufficiently high density. To 
achieve good characteristics of the sintered body besides the high 
strength and high thermal conductivity, the amount of alumina to be 
contained is preferably limited to a range of 0.2-1.5% by weight. 
If alumina is used together with aluminum nitride (AlN) which is mentioned 
later, the total amount of the two compounds is preferably limited to at 
most 2.0% by weight. 
Aluminum nitride plays various roles. For example, it suppresses the 
evaporation of silicon nitride and assists the sintering assistant effect 
of the rare earth element during the sintering process. 
The amount of aluminum nitride is limited to at most 2.0% by weight because 
if it is greater than 2.0% by weight, an excessively large portion of the 
grain boundary phase is formed, or aluminum nitride starts dissolving into 
the silicon nitride and, therefore, reduces the thermal conductivity of 
the sinter. Further, the amount of aluminum nitride contained is 
preferably within a range of 0.3-2.0% by weight because if it is less than 
0.3% by weight (or less than 0.1% by weight in the case where aluminum 
nitride is used together with alumina), the sintered body fails to achieve 
a sufficiently high density. To achieve good characteristics of the sinter 
besides the high strength and high thermal conductivity, the amount of 
aluminum nitride to be contained is preferably set to a range of 0.5-1.5% 
by weight. If aluminum nitride (AlN) is used together with alumina 
(Al.sub.2 O.sub.3), the amount of the aluminum nitride contained in a 
silicon nitride powder is preferably set to a range of 0.1-2.0% by weight. 
The oxides, carbides, nitrides, silicides and borides of Ti, Zr, Hf, V, Nb, 
Ta, Cr, Mo and W promotes the sintering assistant effect of a rare earth 
element, and promotes dispersion thereof in the crystal structure so as to 
enhance the mechanical strength of the silicon nitride (Si.sub.3 N.sub.4) 
sintered body. 
The amount of these compounds contained in a silicon nitride powder is set 
to a range of 0.2-3.0% by weight. If the amount of these compounds 
contained is less than 0.2% by weight, the sintered body fails to achieve 
a sufficiently high density. If the amount is greater than 3.0% by weight, 
the thermal conductivity, mechanical strength and electrical breakdown 
strength of the sintered body are reduced to undesired levels. The 
preferred range of the amount of these compounds contained is 0.3-2.0% by 
weight. 
The above compounds, such as Ti, Zr and Hf, also serve as light blocking 
agents. More specifically, they stain or color the silicon nitride 
sintered body and thus provides it with an opacity. If the silicon nitride 
sintered body of the present invention is used for a base board for an 
integrated circuit or the like which tends to malfunction when exposed to 
light, an appropriate amount of one or more of the above compounds is 
preferably contained in the silicon nitride powder so as to enhance the 
light blocking characteristics of the silicon nitride sintered body. 
Because aluminum compounds, such as aluminum nitride and alumina, assist 
the effect of a sintering assistant agent during the sintering process, 
the use of an aluminum compound relatively reduces the amount of the 
oxides, carbides, nitrides, silicides and borides of Ti, Zr, Hf, V, Nb, 
Ta, Cr, Mo and W needed in a silicon nitride powder. The amount of an 
aluminum compound, such as aluminum nitride or alumina, needed in a 
silicon nitride powder is closely related to the amount of the 
above-mentioned oxides and the like of Ti, Zr, Hf, V, Nb, Ta, Cr, Mo and W 
needed therein. If one or more of aluminum compounds, for example, alumina 
and/or aluminum nitride, are contained in a silicon nitride powder and the 
aluminum compound content is less than 0.1% by weight when the amount of a 
Ti compound or the like as indicated above is less than 0.2% by weight, 
the resulting sintered body fails to achieve a sufficiently high density. 
Further, if the aluminum compound content is greater than 2.0% by weight, 
an excessively large portion of the grain boundary phase is formed, or the 
aluminum compound starts dissolving into the silicon nitride and, 
therefore, reduces the thermal conductivity of the sintered body. 
Therefore, the aluminum compound content must be set to a range of 
0.1-2.0% by weight. To achieve good characteristics of the sintered body 
besides the high strength and high thermal conductivity, the aluminum 
compound content is preferably within a range of 0.2-1.5% by weight. 
The porosity of a sintered body significantly affects the thermal 
conductivity and strength of the sintered body. According to the present 
invention, the porosity is preferably set to 1.5% or less. If it is 
greater than 1.5%, the thermal conductivity and strength of the sintered 
body are reduced to undesired levels. 
Further, the grain boundary phase formed in the silicon nitride crystal 
structure in a sintered body significantly affects the thermal 
conductivity of the sintered body. According to the present invention, the 
ratio of the area of a crystal phase formed in the grain boundary phase to 
the entire area of the grain boundary phase is preferably 20% or greater. 
If the ratio is less than 20%, the sintered body fails to achieve a 
thermal conductivity of at least 60 W/m.multidot.K and, therefore, good 
heat-radiating characteristics and a desirable high-temperature strength. 
To ensure that the porosity of the silicon nitride sintered body is limited 
to at most 1.5% and the ratio of the area of a crystal phase formed in the 
grain boundary phase to the entire area of the grain boundary phase 
becomes at least 20%, a silicon nitride molded compact must be 
pressure-sintered at 1800.degree.-2000.degree. C. for about 0.5-10 hours 
immediately followed by cooling the sintered body at a rate of 100.degree. 
C. per hour or slower. 
If the sintering temperature is lower than 1800.degree. C., the sintered 
body fails to achieve a sufficiently high density; more specifically, the 
porosity becomes greater than 1.5 vol %, thereby reducing both the 
mechanical strength and thermal conductivity of the sintered body to 
undesired levels. If the sintering temperature is higher than 2000.degree. 
C., the silicon nitride per se becomes likely to evaporate or decompose. 
The decomposition and evaporation of the silicon nitride may occur at 
about 1800.degree. C. if the sintering process is performed under the 
normal pressure. 
The rate of cooling a sintered body must be carefully controlled in order 
to achieve crystallization of the grain boundary phase. If the cooling 
rate is faster than 100.degree. C. per hour, the grain boundary phase of 
the structure of the sintered body becomes an amorphous phase (a glass 
phase) and, therefore, the ratio of the area of a crystal phase formed of 
the liquid phase formed during the sintering process to the entire area of 
the grain boundary phase becomes less than 20%. Thereby, the strength and 
thermal conductivity of the sintered body are reduced to undesired levels. 
The sufficiently broad temperature range in which the cooling rate must be 
precisely controlled is from a predetermined sintering temperature 
(1800.degree.-2000.degree. C.) to the solidifying point of the liquid 
phase formed by the reaction of a sintering assistant agent as described 
above. The liquid phase solidifies at about 1600.degree.-1500.degree. C. 
if a sintering assistant agent is used. By maintaining the cooling rate at 
100.degree. C. per hour or slower, preferably, 50.degree. C. per hour or 
slower, at least in a temperature range from the sintering temperature to 
the solidifying point of the liquid phase, most of the grain boundary 
phase becomes a crystal phase, thus achieving a sintered body having both 
high mechanical strength and high thermal conductivity. 
A silicon nitride sintered body according to the present invention can be 
produced by, for example, the following processes. A material mixture is 
prepared by adding predetermined amount of a sintering assistant agent, a 
required additive, such as an organic binder, and alumina, aluminum 
nitride, and/or a compound of Ti, Zr, Hf or the like, to a fine powder of 
silicon nitride which has a predetermined average grain size and contains 
very small amounts of impurities. The material mixture is then molded into 
a compact having a predetermined shape by, for example, a conventional 
sheet molding method, such as the die-press method or the doctor-blade 
method. After the molding process, the molded compact is maintained at 
600.degree.-800.degree. C. for 1-2 hours in a non-oxidizing atmosphere, 
thereby degreasing the compact, that is, thoroughly removing the organic 
binder added in the material mixture preparing process. The degreased 
compact is sintered at 1800.degree.-2000.degree. C. in an atmosphere of an 
inert gas, such as nitrogen gas or argon gas while being pressured by the 
atmosphere gas. 
The silicon nitride sintered body thus produced achieves a porosity of 1.5% 
or less, a thermal conductivity of 60 W/m.multidot.K (25.degree. C.) or 
greater, and a three-point bending strength of 80 kg/mm.sup.2 or greater.

DESCRIPTION OF THE PREFERRED EMBODIMENTS 
The present invention will be further described with reference to Examples 
and Comparative Examples. 
[EXAMPLES 1-3] 
5% by weight of yttrium oxide (Y.sub.2 O.sub.3) powder, that is, a 
sintering assistant agent, having an average grain size of 0.7 .mu.m, and 
1.5% by weight of alumina (Al.sub.2 O.sub.3) powder having an average 
grain size of 0.5 .mu.m were added to a silicon nitride material powder 
having an average grain size of 0.55 .mu.m. The silicon nitride material 
powder used in Examples 1-3 contained 1.3% by weight of oxygen and 0.15% 
by weight of impurity cationic elements with respect to the amount of the 
silicon nitride material powder, and the silicon nitride contained in the 
silicon nitride material powder contained 97% by weight of alpha-phase 
type silicon nitride with respect to the entire amount of the silicon 
nitride. The above-described mixture was wet-blended in ethyl alcohol for 
24 hours and then dried to obtain a material mixture powder. A 
predetermined amount of an organic binder was added to the material 
mixture powder, and then homogeneously mixed. The mixture was then 
press-molded under a molding pressure of 1000 kg/cm.sup.2 to obtain a 
plurality of molded compacts each having a size of 50 mm (length).times.50 
mm (width).times.5 mm (thickness). After the compacts were degreased in 
the atmosphere gas at 700.degree. C. for 2 hours, the compacts were 
sintered to enhance the density thereof, in the following manner. The 
degreased compacts were maintained at 1900.degree. C. for 6 hours in a 
nitrogen gas atmosphere at 7.5 arm to form sintered bodies. While the 
sintered bodies were subsequently cooled, the cooling rates were 
determined at 100.degree. C./hr (Example 1), 50.degree. C./hr (Example 2) 
and 25.degree. C./hr (Example 3) by controlling the power supplied to the 
heating devices provided in the sintering furnaces until the temperature 
inside the furnaces reached 1500.degree. C. Silicon nitride ceramic 
sintered bodies of Examples 1-3 were thus prepared. 
[COMATIVE EXAMPLE 1] 
Silicon nitride sintered body of Comparative Example 1 was prepared in 
generally the same manner as in Example 1, except that after the sintering 
process, the heating device was powered off to cool the sintered body at a 
cooling rate of about 500.degree. C./hr as performed by the conventional 
furnace cooling method. 
COMATIVE EXAMPLE 2] 
Silicon nitride sintered body of Comparative Example 2 was prepared in 
generally the same manner as in Example 1, except for using a silicon 
nitride material powder having an average grain size of 0.60 .mu.m and 
containing 1.5% by weight of oxygen and 0.6% by weight of impurity 
cationic elements, the proportion of alpha-phase type silicon nitride to 
the entire amount of silicon nitride being 93% by weight, instead of the 
silicon nitride material powder used in Example 1. 
[COMATIVE EXAMPLE 3] 
Silicon nitride sintered body of Comparative Example 3 was prepared in 
generally the same manner as in Example 1, except for using a silicon 
nitride material powder having an average grain size of 1.1 .mu.m and 
containing 1.7% by weight of oxygen and 0.7% by weight of impurity 
cationic elements, the proportion of alpha-phase type silicon nitride to 
the entire amount of silicon nitride being 91% by weight, instead of the 
silicon nitride material powder used in Example 1. 
The silicon nitride sintered bodies of each of Examples 1-3 and Comparative 
Examples 1-3 were examined to determine their porosities, thermal 
conductivities, and three-point bending strengths at a room temperature. 
Further, X-ray analysis of each sintered body was performed to determine 
the proportion of the crystal phase to the grain boundary phase (area 
ratio). The obtained values were averaged respectively for Examples or 
Comparative Examples. The results are shown in Table 1. 
TABLE 1 
______________________________________ 
Cooling Thermal 
3-Point 
Rate Crystal 
Conduc- 
Bending 
until Phase tivity Strength 
Sintered 
1500.degree. C. 
Porosity Ratio (W/m .multidot. 
(kg/ 
Body (.degree.C./hr) 
(%) (%) K) mm.sup.2) 
______________________________________ 
Ex. 1 100 0.2 30 70 102 
Ex. 2 50 0.2 50 92 101 
Ex. 3 25 0.2 80 115 98 
C. Ex. 1 
500 0.2 0 40 100 
C. Ex. 2 
100 0.3 0 27 90 
C. Ex. 3 
100 2.5 0 20 78 
______________________________________ 
As indicated in Table 1, the silicon nitride sintered bodies of Examples 
1-3, which were cooled at rates lower than the cooling rate in Comparative 
Example 1 immediately after the density-enhancing sintering process, had 
crystal phases formed in the grain boundary phases. A silicon nitride 
sintered body having a larger proportion of the crystal phase to the grain 
boundary phase achieved a greater thermal conductivity. The three-point 
bending strengths of the sintered bodies of Examples 1-3 were 
substantially the same despite their different crystal phase proportions. 
Thus, the sintered bodies of Examples 1-3 achieved both high strength and 
high thermal conductivity. 
On the other hand, the sintered bodies of Comparative Example 1, which were 
cooled at a high rate of 500.degree. C./hr, formed no crystal phase in the 
grain boundary phase, that is, the entire grain boundary phase was 
amorphous. The thermal conductivity of Comparative Example 1 was 
accordingly low. The sintered bodies of Comparative Example 2, which were 
formed of a silicon nitride material powder containing an increased amount 
of impurity cationic elements, that is, 6% by weight, formed no crystal 
phase in the grain boundary phase although the cooling rate was the same 
as in Example 1. The thermal conductivity of Comparative Example 2 was 
low. The sintered bodies of Comparative Example 3, which were formed of a 
silicon nitride powder having a larger average grain size, that is, 1.1 
.mu.m, had a large porosity indicating an insufficiently low density. As a 
result, the thermal conductivity and strength of Comparative Example 3 
were low. 
[EXAMPLES 4-12 AND COMATIVE EXAMPLES 4-7] 
The material mixture powders of Examples 4-12 were prepared by varying the 
amounts of the same silicon nitride material powder, Y.sub.2 O.sub.3 
powder and Al.sub.2 O.sub.3 powder as used in Example 1, as shown in Table 
2. After the material mixture powders were molded to form compacts and 
degreased the compacts in generally the same manner as in Example 1, the 
degreased compacts were sintered under the conditions as shown in Table 2. 
The silicon nitride ceramic sintered bodies of Examples 4-12 were thus 
prepared. 
The material mixture powders of Comparative Examples 4-7 were respectively 
prepared as indicated in Table 2. More specifically, the Al.sub.2 O.sub.3 
content was significantly reduced in Comparative Example 4. The Y.sub.2 
O.sub.3 content was significantly reduced in Comparative Example 5. The 
Al.sub.2 O.sub.3 content was significantly increased in Comparative 
Example 6. The Y.sub.2 O.sub.3 content was significantly increased in 
Comparative Example 7. 
The material mixture powders were processed in generally the same manner as 
in Example 1, thus obtaining the silicon nitride sintered bodies of 
Comparative Examples 4-7. 
The porosities, thermal conductivities, three-point bending strengths at a 
room temperature (25.degree. C.) and crystal phase-to-grain boundary phase 
proportions (by X-ray analysis) of the sintered bodies of Examples 4-12 
and Comparative Examples 4-7 were determined under the same conditions as 
in Example 1. The results are shown in Table 2. 
TABLE 2 
______________________________________ 
Sintering Cooling 
Conditions Rate until 
Sintered 
Composition (wt %) 
Temp. .times. Time 
1500.degree. C. 
Body Si.sub.3 N.sub.4 
Y.sub.2 O.sub.3 
Al.sub.2 O.sub.3 
(.degree.C.) (hr) 
(.degree.C./hr) 
______________________________________ 
Ex. 4 94.5 5 0.5 1900 .times. 6 
50 
Ex. 5 94 5 1 1900 .times. 6 
50 
Ex. 6 93 5 2 1900 .times. 6 
50 
Ex. 7 96 2 2 1900 .times. 6 
50 
Ex. 8 92 7.5 0.3 1900 .times. 6 
50 
Ex. 9 94 5 1 1900 .times. 6 
100 
Ex. 10 94 5 1 1900 .times. 6 
25 
Ex. 11 97 2 1 1950 .times. 6 
50 
Ex. 12 94 5 1 1950 .times. 6 
10 
C. Ex. 4 
94.8 5 0.1 1900 .times. 6 
100 
C. Ex. 5 
97 1 2 1900 .times. 6 
100 
C. Ex. 6 
92 5 3 1900 .times. 6 
100 
C. Ex. 7 
89 10 1 1900 .times. 6 
100 
______________________________________ 
Crystal Thermal 3-Point 
Phase Conduc- Bending 
Sintered Porosity Ratio tivity Strength 
Body (%) (%) (W/m .multidot. K) 
(kg/mm.sup.2) 
______________________________________ 
Example 4 0.4 55 85 94 
Example 5 0.3 50 88 97 
Example 6 0.2 45 90 102 
Example 7 0.9 25 62 90 
Example 8 0.3 92 120 101 
Example 9 0.2 42 80 100 
Example 10 
0.2 90 120 97 
Example 11 
1.2 45 81 95 
Example 12 
0.1 95 128 96 
Comp. Ex. 4 
2.5 50 51 80 
Comp. Ex. 5 
3.0 15 35 72 
Comp. Ex. 6 
0.1 10 40 105 
Comp. Ex. 7 
0.1 38 50 85 
______________________________________ 
As shown in Table 2, the sintered bodies of Examples 4-12, which contained 
amounts of Y.sub.2 O.sub.3 and Al.sub.2 O.sub.3 within the ranges 
according to the present invention, and which were cooled at predetermined 
cooling rates according to the present invention, achieved sufficiently 
high strengths and thermal conductivities. On the other hand, the sintered 
bodies of Comparative Examples 4-7, in each of which the Y.sub.2 O.sub.3 
content and/or the Al.sub.2 O.sub.3 content was out of the respective 
ranges determined according to the present invention, failed to achieve a 
sufficiently high density or a sufficiently high crystal phase-to-grain 
boundary phase proportion (in a case, an excessively large grain boundary 
phase was formed). Therefore, the bending strength or the thermal 
conductivity thereof was undesirably reduced. 
[EXAMPLES 13-16] 
Silicon nitride ceramic sintered bodies of Example 13-16 were prepared in 
generally the same manner as in Example 1, except that the oxides of rare 
earth elements as shown in Table 3 were used instead of the Y.sub.2 
O.sub.3 powder. 
The porosities, thermal conductivities, three-point bending strengths at a 
room temperature (25.degree. C.) and crystal phase-to-grain boundary phase 
proportions (by X-ray analysis) of the sintered bodies of Examples 13-16 
were determined under the same conditions as in Example 1. The results are 
shown in Table 3. 
TABLE 3 
______________________________________ 
Oxide of Crystal 
Thermal 3-Point 
Rare Phase Conduc- Bending 
Sintered 
Earth Porosity Ratio tivity Strength 
Body Element (%) (%) (W/m .multidot. K) 
(kg/mm.sup.2) 
______________________________________ 
Ex. 13 CeO.sub.2 
0.1 50 88 98 
Ex. 14 Nd.sub.2 O.sub.3 
0.2 55 90 95 
Ex. 15 Yb.sub.2 O.sub.3 
0.2 65 100 102 
Ex. 16 Dy.sub.2 O.sub.3 
0.1 55 98 100 
______________________________________ 
As shown in Table 3, the sintered bodies of Examples 13-16 employing rare 
earth element oxides other than Y.sub.2 O.sub.3 achieved generally the 
same properties as those of the sintered bodies employing Y.sub.2 O.sub.3. 
Sintered bodies employing aluminum nitride (AlN) will be described below. 
[EXAMPLES 17-19] 
5% by weight of yttrium oxide (Y.sub.2 O.sub.3) powder, that is, a 
sintering assistant agent, having an average grain size of 0.7 .mu.m, and 
1% by weight of aluminum nitride (AlN) powder having an average grain size 
of 0.8 .mu.m were added to a silicon nitride material powder having an 
average grain size of 0.55 .mu.m. The silicon nitride material powder used 
in Examples 17-19 contained 1.3% by weight of oxygen and 0.15% by weight 
of impurity cationic elements with respect to the amount of the silicon 
nitride material powder, and the silicon nitride contained in the silicon 
nitride material powder contained 97% by weight of alpha-phase type 
silicon nitride with respect to the entire amount of the silicon nitride. 
The above-described mixture was wet-blended in ethyl alcohol for 24 hours 
and then dried to obtain a material mixture powder. A predetermined amount 
of an organic binder was added to the material mixture powder, and then 
homogeneously mixed. The mixture was then press-molded under a molding 
pressure of 1000 kg/cm.sup.2 to obtain a plurality of compacts each having 
a size of 50 mm (length).times.50 mm (width).times.5 mm (thickness). After 
the compacts were degreased in the atmosphere gas at 700.degree. C. for 2 
hours, the compacts were sintered to enhance the density thereof, in the 
following manner. The degreased compacts were maintained at 1900.degree. 
C. for 6 hours in a nitrogen gas atmosphere at 7.5 atm to form sintered 
bodies. While the sintered bodies were subsequently cooled, the cooling 
rates were determined at 100.degree. C./hr (Example 17), 50.degree. C./hr 
(Example 18) and 25.degree. C./hr (Example 19) by controlling the power 
supplied to the heating devices provided in the sintering furnaces until 
the temperature inside the furnaces reached 1500.degree. C. Silicon 
nitride ceramic sintered bodies of Examples 17-19 were thus prepared. 
[COMATIVE EXAMPLE 8] 
Silicon nitride sintered body of Comparative Example 8 were prepared in 
generally the same manner as in Example 17, except that after the 
sintering process, the heating device was powered off to cool the sintered 
body at a cooling rate of about 500.degree. C./hr as performed by the 
conventional furnace cooling method. 
[COMATIVE EXAMPLE 9] 
Silicon nitride sintered bodies of Comparative Example 9 were prepared in 
generally the same manner as in Example 17, except for using a silicon 
nitride material powder having an average grain size of 0.60 .mu.m and 
containing 1.5% by weight of oxygen and 0.6% by weight of impurity 
cationic elements, the proportion of alpha-phase type silicon nitride to 
the entire amount of silicon nitride being 93% by weight, instead of the 
silicon nitride material powder used in Example 17. 
[COMATIVE EXAMPLE 10] 
Silicon nitride sintered bodies of Comparative Example 10 were prepared in 
generally the same manner as in Example 17, except for using a silicon 
nitride material powder having an average grain size of 1.1 .mu.m and 
containing 1.7% by weight of oxygen and 0.7% by weight of impurity 
cationic elements, the proportion of alpha-phase type silicon nitride to 
the entire amount of silicon nitride being 91% by weight, instead of the 
silicon nitride material powder used in Example 17. 
The silicon nitride sintered bodies of each of Examples 17-19 and 
Comparative Examples 8-10 were examined to determine their porosities, 
thermal conductivities, and three-point bending strengths at a room 
temperature. Further, X-ray analysis of each sintered body was performed 
to determine the proportion of the crystal phase to the grain boundary 
phase (area ratio). The obtained values were averaged respectively for 
Examples or Comparative Examples. The results are shown in Table 4. 
TABLE 4 
______________________________________ 
Cooling 
Rate Crystal 
Thermal 3-Point 
until Phase Conduc- Bending 
Sintered 
1500.degree. C. 
Porosity Ratio tivity Strength 
Body (.degree.C./hr) 
(%) (%) (W/m .multidot. K) 
(kg/mm.sup.2) 
______________________________________ 
Ex. 17 100 0.3 35 72 100 
Ex. 18 50 0.3 55 95 98 
Ex. 19 25 0.2 85 117 93 
C. Ex. 8 
500 0.3 0 48 95 
C. Ex. 9 
100 0.3 0 30 90 
C. Ex. 10 
100 2.8 0 23 70 
______________________________________ 
As indicated in Table 4, the silicon nitride sintered bodies of Examples 
17-19, which were cooled at rates lower than the cooling rate in 
Comparative Example 8 immediately after the density-enhancing sintering 
process, had crystal phases formed in the grain boundary phases. A silicon 
nitride sintered body having a larger crystal phase-to-grain boundary 
phase portion achieved a greater thermal conductivity. The three-point 
bending strengths of the sintered bodies of Examples 17-19 were 
substantially the same despite their different crystal phase proportions. 
Thus, the sintered bodies of Examples 17-19 achieved both high strength 
and high thermal conductivity. 
On the other hand, the sintered bodies of Comparative Example 8, which were 
cooled at a high rate of 500.degree. C./hr, formed no crystal phase in the 
grain boundary phase, that is, the entire grain boundary phase was 
amorphous. The thermal conductivity of Comparative Example 8 was 
accordingly low. The sintered bodies of Comparative Example 9, which were 
formed of a silicon nitride material powder containing an increased amount 
of impurity cationic elements, that is, 6% by weight, formed no crystal 
phase in the grain boundary phase although the cooling rate was the same 
as in Example 17. The thermal conductivity of Comparative Example 9 was 
low. The sintered bodies of Comparative Example 10, which were formed of a 
silicon nitride powder having a larger average grain size, that is, 1.1 
.mu.m, had a large porosity indicating an insufficiently low density. As a 
result, the thermal conductivity and strength of Comparative Example 10 
were low. 
[EXAMPLES 20-31 and COMATIVE EXAMPLES 11-13] 
The material mixture powders of Examples 20-31 were prepared by varying the 
amounts of an Al.sub.2 O.sub.3 powder having an average grain size of 0.5 
.mu.m and the same silicon nitride material powder, Y.sub.2 O.sub.3 powder 
and AlN powder as used in Example 17, as shown in Table 5. After the 
material mixture powders were molded to form compacts and degreased the 
compacts in generally the same manner as in Example 17, the degreased 
compacts were sintered under the conditions as shown in Table 5. The 
silicon nitride ceramic sintered bodies of Examples 20-31 were thus 
prepared. 
The material mixture powders of Comparative Examples 11-13 were 
respectively prepared as indicated in Table 5. More specifically, the 
Y.sub.2 O.sub.3 content was significantly reduced in Comparative Example 
11. The AlN content was significantly increased in Comparative Example 12. 
The Y.sub.2 O.sub.3 content was significantly increased in Comparative 
Example 13. The material mixture powders were processed in generally the 
same manner as in Example 17, thus obtaining the silicon nitride sintered 
bodies of Comparative Examples 11-13. 
The porosities, thermal conductivities, three-point bending strengths at a 
room temperature (25.degree. C.) and crystal phase-to-grain boundary phase 
proportions (by X-ray analysis) of the sintered bodies of Examples (Ex.) 
20-31 and Comparative Examples (C.Ex.) 11-13 were determined under the 
same conditions as in Example 17. The results are shown in Table 5. 
TABLE 5 
______________________________________ 
Sintering Cooling 
Conditions 
Rate 
Temp. .times. 
until 
Sintered 
Composition (wt %) 
Time 1500.degree. C. 
Body Si.sub.3 N.sub.4 
Y.sub.2 O.sub.3 
AlN Al.sub.2 O.sub.3 
(.degree.C.) (hr) 
(.degree.C./hr) 
______________________________________ 
Ex. 20 94.7 5 0.3 1900 .times. 6 
50 
Ex. 21 94 5 1 1900 .times. 6 
50 
Ex. 22 93 5 2 1900 .times. 6 
50 
Ex. 23 96 2 2 1900 .times. 6 
50 
Ex. 24 92.2 7.5 0.3 1900 .times. 6 
50 
Ex. 25 94 5 1 1900 .times. 6 
100 
Ex. 26 94 5 1 1900 .times. 6 
25 
Ex. 27 97 2 1 1950 .times. 6 
50 
Ex. 28 94 5 1 1950 .times. 6 
10 
Ex. 29 94 5 0.5 0.1 1900 .times. 6 
100 
Ex. 30 93.8 5 1 0.2 1900 .times. 6 
100 
Ex. 31 93.7 5 0.3 1 1900 .times. 6 
100 
C. Ex. 11 
97 1 2 1900 .times. 6 
100 
C. Ex. 12 
92 5 3 1900 .times. 6 
100 
C. Ex. 13 
92 10 1 1900 .times. 6 
100 
______________________________________ 
Crystal Thermal 3-Point 
Phase Conduc- Bending 
Sintered Porosity Ratio tivity Strength 
Body (%) (%) (W/m .multidot. K) 
(kg/mm.sup.2) 
______________________________________ 
Example 20 
0.8 48 80 84 
Example 21 
0.4 50 85 95 
Example 22 
0.2 40 80 100 
Example 23 
1.0 30 75 92 
Example 24 
0.5 60 95 98 
Example 25 
0.2 45 80 100 
Example 26 
0.2 80 115 96 
Example 27 
1.3 40 85 90 
Example 28 
0.2 95 120 95 
Example 29 
0.2 50 85 105 
Example 30 
0.2 48 83 108 
Example 31 
0.2 45 81 107 
Comp. Ex. 11 
3.2 20 38 70 
Comp. Ex. 12 
0.1 15 50 100 
Comp. Ex. 13 
0.1 30 55 86 
______________________________________ 
As shown in Table 5, the sintered bodies of Examples 20-31, which contained 
amounts of Y.sub.2 O.sub.3 and AlN and, optionally, Al.sub.2 O.sub.3 
within the ranges according to the present invention, and which were 
cooled at predetermined cooling rates according to the present invention, 
achieved sufficiently high strengths and thermal conductivities. On the 
other hand, the sintered bodies of Comparative Examples 11-13, in each of 
which the Y.sub.2 O.sub.3 content and/or the AlN content was out of the 
respective ranges determined according to the present invention, failed to 
achieve a sufficiently high density or a sufficiently high crystal 
phase-to-grain boundary phase proportion (in a case, an excessively large 
grain boundary phase was formed). Therefore, the bending strength or the 
thermal conductivity thereof was undesirably reduced. 
[EXAMPLES 32-35] 
Silicon nitride ceramic sintered bodies of Example 32-35 were prepared in 
generally the same manner as in Example 17, except that the oxides of rare 
earth elements as shown in Table 6 were used instead of the Y.sub.2 
O.sub.3 powder. The porosities, thermal conductivities, three-point 
bending strengths at a room temperature (25.degree. C.) and crystal 
phase-to-grain boundary phase proportions (by X-ray analysis) of the 
sintered bodies of Examples 32-35 were determined under the same 
conditions as in Example 17. The results are shown in Table 6. 
TABLE 6 
______________________________________ 
Oxide of Crystal 
Thermal 3-Point 
Rare Phase Conduc- Bending 
Sintered 
Earth Porosity Ratio tivity Strength 
Body Element (%) (%) (W/m .multidot. K) 
(kg/mm.sup.2) 
______________________________________ 
Ex. 32 CeO.sub.2 
0.2 55 89 95 
Ex. 33 Nd.sub.2 O.sub.3 
0.3 55 88 98 
Ex. 34 Yb.sub.2 O.sub.3 
0.3 65 103 95 
Ex. 35 Dy.sub.2 O.sub.3 
0.2 60 99 101 
______________________________________ 
As shown in Table 6, the sintered bodies of Examples 32-35 employing rare 
earth element oxides other than Y.sub.2 O.sub.3 achieved generally the 
same properties as those of the sintered bodies employing Y.sub.2 O.sub.3. 
[EXAMPLES 36-38] 
5% by weight of yttrium oxide (Y.sub.2 O.sub.3) powder, that is, a 
sintering assistant agent, having an average grain size of 0.7 .mu.m, and 
1.5% by weight of hafnium oxide (HfO.sub.2) powder having an average grain 
size of 1 .mu.m were added to a silicon nitride material powder having an 
average grain size of 0.55 .mu.m. 
The silicon nitride material powder used in Examples 36-38 contained 1.3% 
by weight of oxygen and 0.15% by weight of impurity cationic elements with 
respect to the amount of the silicon nitride material powder, and the 
silicon nitride contained in the silicon nitride material powder contained 
97% by weight of alpha-phase type silicon nitride with respect to the 
entire amount of the silicon nitride. The above-described mixture was 
wet-blended in ethyl alcohol for 24 hours and then dried to obtain a 
material mixture powder. A predetermined amount of an organic binder was 
added to the material mixture powder, and then homogeneously mixed. The 
mixture was then press-molded under a molding pressure of 1000 kg/cm.sup.2 
to obtain a plurality of molded compacts each having a size of 50 mm 
(length).times.50 mm (width).times.5 mm (thickness). After the compacts 
were degreased in the atmosphere gas at 700.degree. C. for 2 hours, the 
compacts were sintered to enhance the density thereof, in the following 
manner. The degreased compacts were maintained at 1900.degree. C. for 6 
hours in a nitrogen gas atmosphere at 7.5 arm to form sintered bodies. 
While the sintered bodies were subsequently cooled, the cooling rates were 
determined at 100.degree. C./hr (Example 36), 50.degree. C./hr (Example 
37) and 25.degree. C./hr (Example 38) by controlling the power supplied to 
the heating devices provided in the sintering furnaces until the 
temperature inside the furnaces reached 1500.degree. C. Silicon nitride 
ceramic sintered bodies of Examples 36-38 were thus prepared. 
[COMATIVE EXAMPLE 14] 
Silicon nitride sintered bodies of Comparative Example 14 were prepared in 
generally the same manner as in Example 36, except that after the 
sintering process, the heating device was powered off to cool the sintered 
body at a cooling rate of about 500.degree. C./hr as performed by the 
conventional furnace cooling method. 
[COMATIVE EXAMPLE 15] 
Silicon nitride sintered bodies of Comparative Example 15 were prepared in 
generally the same manner as in Example 36, except for using a silicon 
nitride material powder having an average grain size of 0.60 .mu.m and 
containing 1.5% by weight of oxygen and 0.6% by weight of impurity 
cationic elements, the proportion of alpha-phase type silicon nitride to 
the entire amount of silicon nitride being 93% by weight, instead of the 
silicon nitride material powder used in Example 36. 
[COMATIVE EXAMPLE 16] 
Silicon nitride sintered bodies of Comparative Example 16 were prepared in 
generally the same manner as in Example 36, except for using a silicon 
nitride material powder having an average grain size of 1.1 .mu.m and 
containing 1.7% by weight of oxygen and 0.7% by weight of impurity 
cationic elements, the proportion of alpha-phase type silicon nitride to 
the entire amount of silicon nitride being 91% by weight, instead of the 
silicon nitride material powder used in Example 36. 
The silicon nitride sintered bodies of each of Examples 36-38 and 
Comparative Examples 14-16 were examined to determine their porosities, 
thermal conductivities, and three-point bending strengths at a room 
temperature. Further, X-ray analysis of each sintered body was performed 
to determine the proportion of the crystal phase to the grain boundary 
phase (area ratio). The obtained values were averaged respectively for 
Examples or Comparative Examples. The results are shown in Table 7. 
TABLE 7 
______________________________________ 
Cooling 
Rate Crystal 
Thermal 3-Point 
until Phase Conduc- Bending 
Sintered 
1500.degree. C. 
Porosity Ratio tivity Strength 
Body (.degree.C./hr) 
(%) (%) (W/m .multidot. K) 
(kg/mm.sup.2) 
______________________________________ 
Ex. 36 100 0.4 50 89 100 
Ex. 37 50 0.4 70 98 100 
Ex. 38 25 0.3 87 105 98 
C. Ex. 14 
500 0.4 0 48 95 
C. Ex. 15 
100 0.3 0 35 90 
C. Ex. 16 
100 3.0 0 22 70 
______________________________________ 
As indicated in Table 7, the silicon nitride sintered bodies of Examples 
36-38, which were cooled at rates lower than the cooling rate in 
Comparative Example 14 immediately after the density-enhancing sintering 
process, had crystal phases formed in the grain boundary phases. A silicon 
nitride sintered body having a larger crystal phase-to-grain boundary 
phase proportion achieved a greater thermal conductivity. The three-point 
bending strengths of the sintered bodies of Examples 36-38 were 
substantially the same despite their different crystal phase proportions. 
Thus, the sintered bodies of Examples 36-38 achieved both high strength 
and high thermal conductivity. 
On the other hand, the sintered bodies of Comparative Example 14, which 
were cooled at a high rate of 500.degree. C./hr, formed no crystal phase 
in the grain boundary phase, that is, the entire grain boundary phase was 
amorphous. The thermal conductivity of Comparative Example 14 was 
accordingly low. The sintered bodies of Comparative Example 15, which were 
formed of a silicon nitride material powder containing an increased amount 
of impurity cationic elements, that is, 6% by weight, formed no crystal 
phase in the grain boundary phase although the cooling rate was the same 
as in Example 36. The thermal conductivity thereof was low. The sintered 
bodies of Comparative Example 16, which were formed of a silicon nitride 
powder having a larger average grain size, that is, 1.1 .mu.m, had a large 
porosity indicating an insufficiently low density. As a result, the 
thermal conductivity and strength of Comparative Example 16 were low. 
[EXAMPLES 39-69 AND COMATIVE EXAMPLES 17-23] 
The material mixture powders of Examples 39-69 were prepared by varying the 
amounts of the same silicon nitride material powder, Y.sub.2 O.sub.3 
powder and HfO.sub.2 powder as used in Example 36 or other metal compound 
powders and, optionally, Al.sub.2 O.sub.3 powder and/or AlN powder, as 
shown in Tables 8 and 9. After the material mixture powders were molded to 
form compacts and degreased the compacts in generally the same manner as 
in Example 36, the degreased compacts were sintered under the conditions 
as shown in Tables 8, 9 to form sintered bodies. The silicon nitride 
ceramic sintered bodies of Examples 39-69 were thus prepared. 
The material mixture powders of Comparative Examples 17-23 were 
respectively prepared as indicated in Table 9. More specifically, a 
significantly reduced amount of HfO.sub.2 was used in Comparative Example 
17. A significantly reduced amount of Y.sub.2 O.sub.3 was used in 
Comparative Example 18. A significantly increased amount of HfO.sub.2 was 
used in Comparative Example 19. A significantly increased amount of 
Y.sub.2 O.sub.3 was used in Comparative Example 20. A significantly 
increased amount of TiO.sub.2 was used in Comparative Example 21. A 
significantly increased amount of AlN was used in Comparative Example 22. 
A significantly increased amount of alumina was used in Comparative 
Example 23. The material mixture powders were processed in generally the 
same manner as in Example 36, thus obtaining the silicon nitride sintered 
bodies of Comparative Examples 17-23. 
The porosities, thermal conductivities, three-point bending strengths at a 
room temperature (25.degree. C.) and crystal phase-to-grain boundary phase 
proportions (by X-ray analysis) of the sintered bodies of Examples 39-69 
and Comparative Examples 17-23 were determined under the same conditions 
as in Example 36. 
The results are shown in Tables 8, 9. 
TABLE 8 
__________________________________________________________________________ 
Sintering 
Cooling Crystal 
Thermal 
3-Point 
Composition (wt %) Conditions 
Rate until Phase 
Conduc- 
Bending 
Sintered Other Temp. .times. Time 
1500.degree. C. 
Porosity 
Ratio 
tivity 
Strength 
Body Si.sub.3 N.sub.4 
Y.sub.2 O.sub.3 
Component 
Al.sub.2 O.sub.3 
AlN 
(.degree.C.) (hr) 
(.degree.C./hr) 
(%) (%) (W/m .multidot. 
(kg/mm.sup.2) 
__________________________________________________________________________ 
Ex. 39 
94.8 
5 HfO.sub.2 
0.2 1900 .times. 6 
50 1.3 50 88 85 
Ex. 40 
94 5 HfO.sub.2 
1 1900 .times. 6 
50 0.2 55 90 98 
Ex. 41 
92 5 HfO.sub.2 
3 1900 .times. 6 
50 0.1 45 75 108 
Ex. 42 
95 2 HfO.sub.2 
3 1900 .times. 6 
50 0.5 30 65 95 
Ex. 43 
92 7.5 
HfO.sub.2 
0.5 1900 .times. 6 
50 0.1 75 95 88 
Ex. 44 
94 5 HfO.sub.2 
1 1900 .times. 6 
100 0.2 45 80 100 
Ex. 45 
94 5 HfC 1 1900 .times. 6 
100 0.3 55 92 98 
Ex. 46 
94 5 HfN 1 1900 .times. 6 
100 0.4 55 90 95 
Ex. 47 
94 5 HfSi.sub.2 
1 1900 .times. 6 
100 0.4 50 88 93 
Ex. 48 
94 5 HfB.sub.2 
1 1900 .times. 6 
100 0.4 45 83 90 
Ex. 49 
94 5 TiO.sub.2 
1 1900 .times. 6 
100 0.2 48 80 98 
Ex. 50 
95 5 ZrO.sub.2 
1 1900 .times. 6 
100 0.2 45 85 105 
Ex. 51 
94 5 V.sub.2 O.sub.5 
1 1900 .times. 6 
100 0.3 50 90 95 
Ex. 52 
94 5 Nb.sub.2 O.sub.5 
1 1900 .times. 6 
100 0.2 48 83 90 
Ex. 53 
94 5 Ta.sub.2 O.sub.5 
1 1900 .times. 6 
100 0.4 45 80 92 
Ex. 54 
94 5 Cr.sub.2 O.sub.3 
1 1900 .times. 6 
100 0.2 58 100 95 
Ex. 55 
94 5 MoO.sub.3 
1 1900 .times. 6 
100 0.4 40 75 93 
Ex. 56 
94 5 WO.sub.3 
1 1900 .times. 6 
100 0.3 40 75 90 
Ex. 57 
94 5 TiC 1 1900 .times. 6 
100 0.4 59 95 95 
Ex. 58 
94 5 WC 1 1900 .times. 6 
100 0.3 49 83 93 
Ex. 59 
94 5 TiB.sub.2 
1 1900 .times. 6 
100 0.4 40 80 97 
Ex. 60 
94 5 HfO.sub.2 
0.5 1900 .times. 6 
100 0.2 46 82 102 
TiO.sub.2 
0.5 
__________________________________________________________________________ 
TABLE 9 
__________________________________________________________________________ 
Sintering 
Cooling Crystal 
Thermal 
3-Point 
Composition (wt %) Conditions 
Rate until Phase 
Conduc- 
Bending 
Sintered Other Temp. .times. Time 
1500.degree. C. 
Porosity 
Ratio 
tivity 
Strength 
Body Si.sub.3 N.sub.4 
Y.sub.2 O.sub.3 
Component 
Al.sub.2 O.sub.3 
AlN 
(.degree.C.) (hr) 
(.degree.C./hr) 
(%) (%) (W/m .multidot. 
(kg/mm.sup.2) 
__________________________________________________________________________ 
Ex. 61 
94 5 ZrO.sub.2 
0.5 1900 .times. 6 
100 0.2 52 84 98 
Cr.sub.2 O.sub.3 
0.5 
Ex. 62 
94 5 TiC 0.5 1900 .times. 6 
100 0.2 50 90 100 
HfO.sub.2 
0.5 
Ex. 63 
94.4 
5 HfO.sub.2 
0.5 
0.1 1900 .times. 6 
100 0.3 50 85 98 
Ex. 64 
92.8 
5 TiC 0.2 
2 1900 .times. 6 
100 0.2 30 70 100 
Ex. 65 
94 5 HfO.sub.2 
0.5 0.5 
1900 .times. 6 
100 0.1 58 90 99 
Ex. 66 
92.8 
5 HfO.sub.2 
0.2 2 1900 .times. 6 
100 0.2 40 75 90 
Ex. 67 
94.1 
5 HfO.sub.2 
0.5 
0.1 0.3 
1900 .times. 6 
100 0.2 60 98 102 
Ex. 68 
92.8 
5 Cr.sub.2 O.sub.3 
0.2 
1 1 1900 .times. 6 
100 0.2 40 68 105 
Ex. 69 
92.8 
5 TiO.sub.2 
2 0.1 0.1 
1900 .times. 6 
100 0.2 45 76 100 
C. Ex. 17 
95.9 
5 HfO.sub.2 
0.1 1900 .times. 6 
100 2.6 55 60 78 
C. Ex. 18 
97 1 HfO.sub.2 
2 1900 .times. 6 
100 5.5 20 25 68 
C. Ex. 19 
90 5 HfO.sub.2 
5 1900 .times. 6 
100 0.1 18 50 88 
C. Ex. 20 
89.5 
10 HfO.sub.2 
0.5 1900 .times. 6 
100 0.2 20 52 85 
C. Ex. 21 
90 5 TiO.sub.2 
5 1900 .times. 6 
100 0.1 15 48 90 
C. Ex. 22 
91.5 
5 HfO.sub.2 
0.5 3 1900 .times. 6 
100 0.2 14 50 80 
C. Ex. 23 
91.9 
5 TiO.sub.2 
0.2 
3 1900 .times. 6 
100 0.2 5 40 98 
__________________________________________________________________________ 
As shown in Tables 8 and 9, the sintered bodies of Examples 39-69, which 
contained amounts of metal compounds, such as Y.sub.2 O.sub.3 and 
HfO.sub.2, and, optionally, Al.sub.2 O.sub.3 and/or AlN within the ranges 
according to the present invention, and which were cooled at predetermined 
cooling rates according to the present invention, achieved sufficiently 
high strengths and thermal conductivities. On the other hand, the sintered 
bodies of Comparative Examples 17-23, each of which contained an 
excessively reduced or increased amount of at least one of Y.sub.2 
O.sub.3, HfO.sub.2, TiO.sub.2, Al.sub.2 O.sub.3 and AlN, that is, an 
amount out of the respective ranges determined according to the present 
invention, failed to achieve a sufficiently high density or a sufficiently 
high crystal phase-to-grain boundary phase proportion (in a case, an 
excessively large grain boundary phase was formed). Therefore, the bending 
strength or the thermal conductivity thereof was undesirably reduced. 
[EXAMPLES 70-73] 
Silicon nitride ceramic sintered bodies of Example 70-73 were prepared in 
generally the same manner as in Example 36, except that the oxides of rare 
earth elements as shown in Table 10 were used instead of the Y.sub.2 
O.sub.3 powder. 
The porosities, thermal conductivities, three-point bending strengths at a 
room temperature (25.degree. C.) and crystal phase-to-grain boundary phase 
proportions (by X-ray analysis) of the sintered bodies of Examples 70-73 
were determined under the same conditions as in Example 36. The results 
are shown in Table 10. 
TABLE 10 
______________________________________ 
Oxide of Crystal 
Thermal 3-Point 
Rare Phase Conduc- Bending 
Sintered 
Earth Porosity Ratio tivity Strength 
Body Element (%) (%) (W/m .multidot. K) 
(kg/mm.sup.2 ) 
______________________________________ 
Ex. 70 CeO.sub.2 
0.3 48 85 95 
Ex. 71 Nd.sub.2 O.sub.3 
0.4 47 85 95 
Ex. 72 Yb.sub.2 O.sub.3 
0.5 55 87 90 
Ex. 73 Dy.sub.2 O.sub.3 
0.4 50 87 98 
______________________________________ 
As shown in Table 10, the sintered bodies of Examples 70-73 employing rare 
earth element oxides other than Y.sub.2 O.sub.3 achieved generally the 
same properties as those of the sintered bodies employing Y.sub.2 O.sub.3. 
In addition to the above examples, various material mixture powders were 
prepared by adding 5% by weight of Y.sub.2 O.sub.3 powder and 1 by weight 
of at least one compound selected from the group consisting of ZrC, VC, 
NbC, TaC, Cr.sub.3 C.sub.2, Mo.sub.2 C, TiN, ZrN, VN, TaN, CrN, Mo.sub.2 
N, W.sub.2 N, TiSi.sub.2, ZrSi.sub.2, VSi.sub.2, NbSi.sub.2, TaSi.sub.2, 
CrSi.sub.2, MoSi.sub.2, WSi.sub.2, ZrB.sub.2, VB.sub.2, NbB.sub.2, 
TaB.sub.2, CrB.sub.2, MoB.sub.2 and WB.sub.2, to the silicon nitride 
powder. The material mixture powders were then processed in generally the 
same manner as in Example 36, thus obtaining various silicon nitride 
sintered bodies. The porosities, thermal conductivities, three-point 
bending strengths at a room temperature (25.degree. C.) and crystal 
phase-to-grain boundary phase proportions (by X-ray analysis) of the 
silicon nitride sintered bodies were determined under the same conditions 
as in Example 36. The results were substantially the same as in Examples 
36-73. 
As described above, according to the present invention, a silicon nitride 
sintered body is produced by: molding and sintering a fine powder of 
silicon nitride having predetermined purity and grain size and containing 
predetermined amounts of a rare earth element and, optionally, aluminum 
nitride and/or alumina, or a compound of Ti, Zr, Hf, etc.; and cooling the 
sintered body at a low cooling rate, that is, 100.degree. C. per hour or 
lower. Unlike the conventional process in which sintered body is cooled 
rather quickly by, for example, the furnace cooling method, the grain 
boundary phase of the sintered body is changed from the amorphous state to 
such a state where a crystal phase is present, according to the present 
invention. Thus, the silicon nitride sintered body of the present 
invention achieves high density, high strength and high thermal 
conductivity. The sintered body of the present invention is suitable for 
electronic components and parts, such as semiconductor substrates, 
heat-radiating plates, etc.