Processing and Properties of Advanced Structural Ceramics

5

Processing and Properties of Advanced Structural Ceramics LUDWIG J. GAUCKLER ETH Zürich, CH-8092 Zürich, Switzerland

INTRODUCTION

Advanced structural ceramics are a relatively new class of materials with outstanding properties which have already found a wide range of applications. A few years ago the public became increasingly aware of their possibilities because of the success of the Japanese in pushing these materials forward for many uses. From this success originated the introduction of a classification of the high performance ceramics in terms of their function, properties and applications (Yano, 1984) as shown in Fig. 1. The wide spectrum of uses for these materials ranges through optical, chemical, biological and electrical functions, which make use of their magnetic, mechanical and thermal properties. It continues through to technical applications in the nuclear industries. Advanced structural ceramics can be found in every segment of this wider range of materials. They perform mainly mechanical functions in which properties such as high strength, wear resistance and corrosion resistance are major requirements. The most important examples of advanced structural ceramics are materials based on aluminium oxide, zirconium oxide, silicon nitride and silicon carbide. Usually these materials are not employed in their pure form, but are made into alloys, solid solutions and compounds. Further examples of special advanced structural ceramics are cordierite (MAS), spodumene (LAS), mullite, aluminium titanate and fibre-, particle- and platelet-reinforced ceramics. HIGH-TECH CERAMICS: VIEWPOINTS AND PERSPECTIVES ISBN 0-12-421950-0

Copyright © 1989 Academic Press Limited All rights of reproduction in any form reserved.

L J. Gauckler

60

High temperature| ndustrial furnace lining Electrode material Heat sinl< for electronic^ Tools and jigs Darts,

Nuclear fuel /Nuclear fuel cladding^Control material ^Moderating material Reactor lining > Radiation resistance Laser diodeN / Refractoriness Light-emitting\ High temperature (fi od e strength vHeat-resistant\ /Optical translucent ^ condensing \ p o r c e l a i n Fluorescence \Optical 1 ^ / Nuclear Translucence \ communie. Optical \ cable .conductivity Optical

zi 2

Abrasives Turbine blade /High strengtl / W e a r resistane Thermal / L o w thermal expansior Lubrication, Mechanical Artificial High bone and Biological Performance tooth compatibility/ Ceramics Catalyst Adsorption Biological carrier f

Heat Catalysis 1 exchanger Corrosion Chemica resistance equipment'

Fig. 1.

Chemical

lEIectrical (insulation Electric | Electrical Magnetic I conductivity /Pie/oelectric Dielectric

[IC substrate' Resistance heating element Varistor Sensor Memory element

High performance ceramics: classification according to Yano (1984).

Advanced structural ceramics are manufactured from chemically processed materials and not from natural raw materials as is the case in classical ceramics. The possibility of choosing from many chemical raw materials results in a large number of combinations and compositions. This offers great flexibility in preparing special alloys and compounds starting from an atomistic and microstructural level. A significant and somewhat unfortunate attribute of advanced structural ceramics is their higher cost in comparison to classical ceramics and to most metals. In the following sections the important properties of structural ceramics are briefly reviewed. Special emphasis is given to the technical challenges which are presented to the materials scientist and engineer by the brittle nature of these materials. The various strategies available to improve reliability through developments in processing and examples of microstructu­ ral design to increase toughness and creep resistance are presented. PROPERTIES AND APPLICATIONS

In recent years there has been considerable progress in developing structural ceramics of higher strength and toughness. At present there seems to be no

Processing and Properties of Advanced Structural Ceramics

61

limit to such improvements, except the theoretical bond strength and an upper level as a consequence of the polycrystalline nature of the material. The first is estimated to be between 10 and 40GPa, which is considerably higher than conceivable for most metals. Figure 2 shows schematically the development of materials' strength during recent decades. Besides high mechanical strength, advanced structural ceramics exhibit high melting points because of their strong ionic and covalent bonds. As a consequence they also possess great hardness and wear resistance. They show great stiffness (i.e. high Young's modulus), and their strength is particularly impressive under compressive loading at higher temperatures. Table 1 summarizes the properties of advanced structural ceramics. The high corrosion resistance and oxidation resistance are also very significant. An important contrast to metals is provided by the very low specific weight of these materials. This property together with high strength leads to a very high strength-to-weight ratio which is particularly important for rapidly moving components in machines. Most attractive are the mechanical properties at elevated temperatures. It should be stressed here that no metal can be used under mechanical load at temperatures above 1100°C for any

3,

Fig. 2.

y

r-,^7

¥
Strength development of some materials.

,

ΓΤΤΤΤΊ

* Decomposition temperature.

2300-2500 1750-1900 1750-1900 (1860)

3.95 5.90 6.05 1.5-2.0 3.17 2.6 3.2 3.2

300-370 200 200 130 420 150 230 20

(MPa)

(gem"3)

(°C)

2050 2500-2600 2500-2600

Modu ilus of elasticity

Density

irai ceramics.

Melting point

Properties of several advanced stru

Alumina Zirconia TZP Zirconia PSZ Cordierite (porous) Silicon carbide* Silicon nitride (RBSN)* Silicon nitride (dense)* Aluminium titanate*

Table 1.

3-4.5 6-15 6-12 1.5-2.5 4 2.5-3.5 6

(MPa m*)

Fracture toughness

(GPa) 16-17 12 11 8 20 5-7 15

350-580 900-1000 900-1000 100-200 450-520 200-350 500-600 40

Vickers hardness

(MPa)

Bond strength

8.1 10.5 10.5 0.8-1.2 4.3 3.0 3.3 2.0

Coefficient of thermal expansion (10~ 6 K- 1 )

Processing and Properties of Advanced Structural Ceramics

63

length of time without the use of some sort of cooling system. For these higher temperature applications advanced structural ceramics are the only possible materials, as shown in Fig. 3. An unfortunately important characteristic of these ceramic materials is their brittleness. Ceramics break without any warning if they are subjected to any substantial deformation. They do not demonstrate the plastic characteristics which we normally associate with metals. This property is due to the high ionic and covalent bonding forces as well as to the more complex crystal structures of ceramic materials compared with metals. The low toughness causes problems for designers who are generally used to combining materials which are capable of accepting a certain amount of plastic deformation. Applications for advanced structural ceramics can be divided into two main groups (Table 2). The first list in the table gives applications which have already been realized. This consists basically of components with properties such as wear resistance or corrosion resistance, which are used at lower or intermediate temperature ranges, and in which only moderate mechanical properties are a requirement. Although the range of applications is already very wide, most of the more sophisticated ceramics have only been developed in the last 15 years. For example, the use of ceramic cutting tools has increased threefold since 1972. In the next five years it is anticipated that there will be a further increase by a factor of about three. Use of wear-resistant parts is projected to double within the next five years (Little, 1986). Yet a further example is the anticipated major increase in the catalyst substrates in automotive applications throughout Europe because of recent legislation.

Superalloy IN 100 (cast)

500 Fig. 3.

1000 Temperature (°C)

Bend strength of ceramics and metals.

1500

L J. Gauckler

64 Table 2. Applications of advanced structural ceramics Bearings Medical implants Burners Membranes Casting dies Nozzles Catalyst supports Port liners Coil forms Protection tubes Cutting tools Pump parts Drawing dies Rolling jigs Extrusion dies Textile machine parts Filters Thermal insulators Heaters Turbo charger wheels Laboratory equipment Wear-resistant parts Potential applications Engine components Bearings Cylinder liners Diesel particulate traps Glow plugs Manifolds Piston caps and rings Prechamber Tappets Valves and seats

Gas turbine parts Bearings Blades Heat shields Heat exchanger Housing Rotor and shaft Shrouds Stators

This last example shows one of the major tendencies, namely the increasing requirement for performance at higher temperatures. The second list in Table 2 consists of potential applications. The major potential for structural ceramics is their use as automotive engine components, for gas turbines and energy conversion systems. These applications are characterized by the fact that the ceramic component is an integral part of the machine, and in the event of failure the whole system would cease to function and very frequently could be destroyed. This, therefore, introduces the major requirements of reliability and reproducibility of properties, since the reliability of the ceramic component determines the reliability of the complete system. Various market studies have stated that a large increase in the use of ceramic materials is anticipated (Cadotte et a/., 1987). There are big differences in the projected figures which are related to various uncertainties in relation to the time-scale required for the technical breakthrough to take place. For applications that have already been introduced, the market studies are consistent and suggest a turnover of about US $5 billion (109) for all advanced ceramic materials worldwide. Advanced structural ceramics account for 10% of this total. Cautious estimates project a 10 to 15% annual growth rate for

Processing and Properties of Advanced Structural Ceramics

65

non-automotive applications. However, the largest potential is doubtless presented by the automotive market, and here estimates vary from $7 to $20 billion worldwide. This is a very attractive potential market to stimulate research and development work. TECHNICAL CHALLENGES

The main obstacles which stand in the way of a wider use of structural ceramics are their relatively high cost, their lack of reliability as mechanical components under static and more significantly under dynamic load, and the general lack of knowledge in designing with brittle materials. Costs and cost structures

The most obvious barrier against a wider use of ceramic components is their cost when compared on a one-to-one basis with the same component made of metal. The cost results from high investments in specialized equipment, high labour charges in processing small to medium size lots, and the expensive finishing of ceramic components. These high costs are not only caused by small production runs, but they are also related in part to the specialized properties of ceramics and to difficulties in their manufacture. An example was presented recently of a cost structure analysis of a simple silicon carbide disc (Rothman and Bo wen, 1986). The cost calculation (Fig. 4) showed that the finishing of the component was responsible for 45% of the total cost. In addition, yield of satisfactory components was only 40%, and since most of the defects could only be detected after finishing, scrap was the single greatest cost factor. In fact, the cost of powder was $2.90/kg, whereas the total cost of the component was $187/kg. The challenge is therefore to reduce total manufacturing costs. One major method of reducing costs would be to use better powders. The use of a better Table 3.

Manufacturing costs of a range of structural materials.

US$/kg Main cost items

Advanced Metals structural ceramics

Electronic substrates and ferrites

Earthenware tiles

Structural clay ceramics

50-1000 Labour, capital

1-100 Capital, labour

0.5-5 Labour, raw materials

0.1-0.5 Energy, transportation

5-50 Energy, capital

66

L J. Cauckler [1111111111 Inspection Green Mach

(a)

| ^ | [

Firing | Drying

Powder cost ■ $ 2 . 9 0 / k g Yield : 40% Total cost= $187/kg

\i;fJU

Isopress

WM:^ Material

(b)

|

] Finishing

Powder cost = $ 2 2 / k g Yield ■ 86% Total cost = $122/kg

Fig. 4. Cost distribution of silicon carbide seals (according to Rothman and Bowen, 1986): (a) with low-quality low-cost powder, (b) with high-quality high-cost powder.

and therefore more expensive powder, costing $22/kg, reduced the scrap after every processing step and achieved a total scrap level of less than 14%. The total manufacturing cost was therefore reduced to $122/kg. This example illustrates the importance of a high-quality starting powder. Use of high-quality powders and near-net-shape forming techniques, which reduce finishing costs to a minimum, are two of the methods which are the subject of major development efforts.

Reliability

An important technical problem is the low reliability of these materials in terms of toughness and strength. What is significant here is not the average value of strength of these materials, which is impressive, but rather the distribution of values around a mean, which are unwanted. This is illustrated in Figure 5. The figure shows that the strength of today's commercial sintered silicon carbide (SSiC) lies at an average value of 400 MPa with a Weibull modulus of m = 10. Thereby the strength distribution around this mean ranges from a lower limit of 200 up to 550 MPa. In contrast, the aluminium pressure die cast alloy (G-Al Si 10) shows an average strength of 270 MPa, but this value is true within the tight tolerance range of + 10 MPa. In order to make use of the full potential of the high mechanical strength properties of sintered silicon carbide, it is essential that the scatter of properties be reduced. If the scatter is described in terms of the Weibull statistics parameter m, a measure

Processing and Properties of Advanced Structural Ceramics

67

1.0 h

| Cast Al S±10MPa

0.8 p

£

0.6

Φ 3

σ ω £

SSiC

h

0.4 l·

m = 10LA m = 20

0.2 r

m = 5/\ I L*rS

|

Γν\

|

|

|

100 200 300 400 500 600 700 Fracture Strength (MPa) Fig. 5. Fracture strength distributions of cast aluminium (10wt% Si) and sintered silicon carbide (SSiC) materials (m = Weibull modulus).

of the width of the distribution curve, then a value of m = 10 is typical of a silicon carbide structural ceramic in use today. It is important that this value be improved to over m = 20 to make full use of the potentialities of this material. What then is the reason for this lack of reliability? How can materials scientists arrive at definable quantities for safety factors which are a basic requirement for constructive designers? These questions will be addressed in the following paragraphs. Flaws - the origin of failure in brittle materials

In contrast to metals and plastics, ceramics, at their service temperature, cannot reduce internal and external stresses by the mechanism of plastic deformation. Because of their brittle nature, they are extremely susceptible to cracks and defects both on their surface and in their internal structure. Such defects, however, are almost always present in polycrystalline ceramics. They occur during the manufacturing process or later in service. The

L J. Gauckler

68

approach to the understanding of brittle materials was provided by Griffith in 1920, who showed that when a flaw in the internal structure was greater than a critical value c, then the release of elastic energy was greater than the gain in surface energy. Under these circumstances, the flaw will grow at a catastrophic speed because the driving energy at the flaw tip will increase continuously. The stress at failure can be expressed as follows: -

^

(1)

or

in which E is the elastic modulus, y is the surface energy, c is the length of a surface flaw (2c for volume flaws) and Kc is the critical stress intensity factor, y is a dimensionless geometrical constant. Therefore the largest size of the flaws and their statistical distribution determine the mechanical strength of the material. A major difficulty results from the fact that it is normally impossible to determine the size of the largest defect without destroying the component. However, it is possible to get an idea of the flaw size distribution by tensile and bend tests. The distribution of rupture strength data is the inverse to the flaw size distribution. The third parameter in Equation (2) is Kc, the critical stress intensity factor. Kc is a measure of how tough or how brittle a material is, and it also determines the sensitivity with which the material will react to impact loading. Materials with high Kc values (in the range 10-15 MPam*), are much less sensitive to point loading and impact than materials with Kc values in the range 2-5 M Pam*. Techniques for improving reliability

From the Griffith relation (2) it is possible to illustrate graphically the techniques which are used to improve reliability. Figure 6 shows the strength of present and future ceramics (Abe, 1985). The first and most important step for improving the strength and reliability of monolithic ceramics is to reduce the size of flaws. This can be achieved by obtaining a fine-grained more uniform structure by use of improved powders (more homogeneous, purer) and by adopting better processing techniques (Evans, 1982). The number of defective components can also be

Processing and Properties of Advanced Structural Ceramics

Fig. 6.

69

Strength of present and future ceramics (according to Abe, 1985).

reduced by using non-destructive testing after each processing step. Finally the components can be subjected to proof-testing which will eliminate the weakest components and thereby increase reliability of the remaining components. The second method involves techniques to improve toughness. In brittle materials the critical flaw size is much smaller than in materials which are more ductile. By increasing their toughness, the components become more tolerant in terms of critical flaw size. This method is not only important in terms of making a component more tolerant to defects of the microstructure, but it also improves the reliability in service of the component. A part which is defect-free, but made out of a more brittle material, however, is more likely to fail in service if subjected to shock loading by impacting particles or other defect-introducing events from outside. A further technical challenge is presented by the need to improve the service life of structural ceramics. Even if it was possible to eliminate the larger structural defects, failure can still occur under permanent loading as a result of subcriticai crack growth. Since most cracks start at the surface at grain boundaries, grain boundary engineering can be used to make materials more resistant to slow crack

70

L J. Cauckler

growth. In materials which are intended for use at high temperatures this is particularly important.

Design with brittle materials

Another major barrier to the wider use of advanced structural ceramics is the general lack of knowledge in engineering circles as to how to design with brittle materials. Very frequently the attempt is made to replace a metallic component in a machine with a similarly shaped part made of ceramics. Failure is therefore preprogrammed. When designing components with ceramic materials, full consideration must be given to the shaping possibilities of the ceramic manufacturing process itself and the consequent effect that this process will have on the properties which can be achieved in the component. Methods for design in ceramics were part of the development programme for gas turbines in the United States (Burke et a/., 1978) and Japan (Somiya et a/., 1983). Fig. 7 shows the design concept for brittle materials adopted from E. Munz (personal communication, 1988). This is dependent on the loading of the component, the component geometry as well as the choice of the material. The main tensions for each volume element of a component should be calculated, for example by use of the finite element method. By use of the local failure criteria and statistical property data of the material (σ, m, σ0, ί, etc.), it is possible to calculate the failure probability of each element and of a component. If these calculations lead to an unsatisfac­ tory result, then the material, the component geometry or the loading of the component can be changed until an acceptable solution is found. For this type of analytical procedure, a designer needs the time dependence of the mechanical properties as well as crack propagation velocities, creep rates and cyclic failure data (Hartstock and McLean, 1984). In most cases, however, these data are either not available or available in an unsuitable form. It is important therefore to continue to collect material data, and at the same time to gain more experience with design in ceramics. Design in ceramics has much to offer for the future.

PROCESSING TECHNOLOGIES: STATUS AND OUTLOOK

The main thrust of current research is directed towards improving the classical powder metallurgical manufacturing processes by which most structural ceramics are produced (Harmer et a/., 1986). Some efforts are also

Processing and Properties

of Advanced

Structural

71

Ceramics

Geometry of

Material's

Load on

components

physical properties

component

Main stresses

Mechanical properties

Local failure

m, σ0, t t , v., etc

criterion

Failure probability of c o m p o n e n t

Failure density as function of location and time

Design optimization Fig. 7.

Design concept w i t h brittle materials (E. Munz, personal c o m m u n i c a t i o n , 1988).

being directed towards developing new production techniques. This is evident from the many conferences and publications on this theme (Chen et a/., 1986; Brinker et al.9 1986; Messing et al, 1988). The central objective of all this development effort is to produce ceramics which are either defect-free or at least have a very low level of small defects.

Classical ceramic processing routes

The classical manufacturing process of ceramics begins with powder synthesis. The usual methods are solid reactions or wet chemical reactions with associated calcination of the hydroxide. A schematic diagram of the sequence is shown in Fig. 8. After powder synthesis the powder is prepared, milled, mixed with doping

L J.

72

Powder

Powder

Synthesis

Preparation

h*-i

Forming

Densification

Gauckler

Others

Solid/Solid Reaction

Mixing

Dry Pressing

Sintering

Melting a n d Comminuting

Deagglomeration

Casting

HIP = Hot Isostatic Pressing

Precipitation

Spray D r y i n g

Extrusion

GPS = Gas Pressure Sintering

Freeze D r y i n g

Injection M o l d i n g

HP

Decomposition Gas/Gas Reaction Fig. 8.

= Hot Pressing

Impregnation

Ceramic processing sequence.

agents and with auxiliary organic materials, and then homogenized. In recent years it has been recognized that powder synthesis is an essential element in the production of defect-free ceramics. Uncontrolled agglomerates and impurities at the powder synthesis stage lead to unreliability of the finished product. For the forming process there are a series of techniques available. These range from dry pressing, extruding, tape casting and injection moulding to slip casting. In fact the geometry of the part, the numbers to be produced and the costs as well as specific quality requirements are the determining factors when choosing the process. The next production step is sintering. This has been the focus of development efforts in the last two decades. Sintering is a complex process in which a competition occurs between those reactions causing densification and those causing coarsening of the grains. The aim is to end up with a fine­ grained dense ceramic microstructure. The effect of impurities and dopants on solid- and liquid-phase sintering is relatively well understood when compared with the understanding of the effect of heterogenities in the powder compact. Sintering is made more rapid by use of fine powders and higher temperat­ ures as well as by applying pressure. After sintering it is frequently necessary to finish a component either by sizing to tighter tolerances or by improving the surface finish. At each step of the process it is possible to introduce defects. The individual defects being introduced at each stage differ both in terms of their frequency and their size. Fig. 9 illustrates the size and frequency effect (Lange, 1986; Greil, 1988). A discussion of the individual steps is given below. Examples are also given of the various processes which are most suited to producing defectfree ceramics or to reducing the frequency of such defects.

73

Processing and Properties of Advanced Structural Ceramics

1.0 0.8 £ 0.6 c ω er α>

0.4 0.2

0 1

10

100

1000 Defect Size (μηι)

Fig. 9. Relative incidence of defects introduced in ceramics during manufacturing (according to Lange, 1986; Greil, 1988).

Ceramic powder synthesis

The requirements for a starting powder for a structural ceramic are manysided. One of the most important aspects concerns the purity of the powder in terms of foreign particles (Francis et a/., 1972) as well as the purity of the desired process chemicals themselves, e.g. Na 2 0 in A1203 or Si0 2 in Zr0 2 powders. In addition, the powders should not have hard agglomerates. These agglomerates are one of the causes of defects and play a major part in reducing the strength of components. It is universally accepted that uniform and spherical particles should be used in order to reduce sintering tempera­ tures, thereby avoiding exaggerated grain growth. There has been much work on the synthesis of perfect, monosized powder particles, mainly using a controlled sol-gel technique (Barringer and Bo wen, 1982). Fig. 10 shows such a powder. However, more reliable ceramics were not obtained until efforts were made to eliminate agglomerates in the powder and in the powder compact. By properly controlling the pH of the powder dispersion and by slow sedimentation, uniformally packed green compacts of high green density (about 65% of the theoretical density) could be formed. These compacts sintered at much lower temperatures than powder compacts made from agglomerated powders and resulted infine-graineduniform microstructures. Monosized powders made by hydrolysis of metal alkoxides have particle sizes in the range 0.1-1 ^m. Those made by coprecipitation methods (e.g. Zr0 2 /3mol% Y 2 0 3 ) are in the range 0.02-0.1 μπι.

L J. Gauckler

Fig. 10. Monodispersed borosilicate powder (Courtesy of J. W . Halloran, CPS Supercon­ ductor Corp., Cambridge, Massachusetts).

Alumina

Alumina powder for structural ceramics is manufactured using one of four processes: the Bayer process, thermal decomposition of ammonium-alum, hydrolysis of aluminium alkoxides, and treatment of aluminium flakes by electric power discharge in water (Iwatani process). These four processes are shown schematically in Fig. 11. The Bayer process produces a wide range of aluminas, some of which Bauxite + NaOH

1

Digestion

1

Precipitation

1

AI(OH)3

1

Calcination

AI(OH)3 + H 2 S0 4

AI - Metal

(NHJ2SO4

+

C 3 H 7 OH

NH 4 AI(S0 4 ) 2 12H 2 0

AI(OC 3 H ? ) 3

1 I

Purification

I

Decomposition

I

(x - Al 2 0 3

1 I

Hydrolysis

I

ΛΙ 2 0 3 - Gol

I

Calcination

Al - Metal - Pellets

t

Arc Discharge

| ΛΙ(ΟΗ)3

I

Calcination

I

α-ΛΙ203

I

1

α-ΛΙ20,


Al - Isopropoxide Route

Iwatani - Process

Processing and Properties of Advanced Structural Ceramics

75

have N a 2 0 levels well below 0.01 wt% and varying grain sizes corresponding to specific surface areas of 0.5 up to 2 0 m 2 g _ 1 BET. Aluminas with lower levels of N a 2 0 are produced via the alum process, the Iwatani process or by the hydrolysis of aluminium alkoxides. Fig. 12 shows the N a 2 0 content of commercially available powders plotted against their fineness (in BET surface units). In addition to the fineness of the powder and the N a 2 0 impurity level, Si0 2 , F e 2 0 3 , the T i 0 2 content, and the composition of the various phases (a/y A1 2 0 3 ) as well as the agglomerate factor of the powder are important. These determine to a very large extent the quality of the end-product and are also reflected in the cost of the powders. Deagglomerated powders of high purity with a N a 2 0 content below 0.01 wt% and a high α-Α1 2 0 3 content are used for advanced structural ceramics. A crucial advantage of the manufacturing methods such as the alum or the alkoxide process is the possibility of producing much finer powder without grinding. Usually it is sufficient to deagglomerate these powders and thus obtain the necessary fine dispersion. Using the hydrolysis of aluminium alkoxide, Fanelli and Burlew (1986) were able to produce small particles of alumina powder of about 20 nm in diameter.

^

\

0.1

à*

J

\ \

Bayer - Process Alumina

I I I

O ™ z

v_

0.01

| 0.001 0.1

1

(—Ί I \lw/ Alum - \ / Process

(

\l

10

)

Ì

100 2

Spec. Surface BET (m /g) Fig. 12. Na 2 0 impurity level and specific surface area of Al 2 0 3 powders. Iw, Iwatani process.

L J. Gauckler

76 Zirconia

Zirconia powders for tetragonal zirconia polycrystals (TZP) and partially stabilized zirconia (PSZ) ceramics are also manufactured using wet chemical techniques involving coprecipitation of zirconium salts and other metal salts or hydrolysis of zirconium alkoxides. PSZ powder is also obtained by using a high-temperature melting process and milling. In Fig. 13 these techniques are schematically shown. By use of wet chemical precipitation techniques it is possible to obtain almost monosized powders with particle sizes in the nanometer range. Their shape is nearly spherical after calcination, see Fig. 14. Another method of producing powder is by means of high-temperature melting of zirconia together with its stabilizing oxides, then solidifying from the liquid phase and subsequently milling the resulting material down to the required particle size. With this technique, however, there is the danger that during solidification segregation will occur, making the resulting powder inhomogeneous. Alternative wet chemical techniques, involving deagglomeration of inter­ mediate and end-products, have been discussed by Van de Graaf and Burggraaf (1983). Flame hydrolysis and hydrothermal synthesis have been examined as possible methods of manufacturing ceramic powders by Hirano (1987) and Kriechbaum et al. (1988). Powders have been produced using the hydroxide gel method and these have the advantage of showing very little tendency to build agglomerates (Duran et ai, 1988). Other methods are being investigated, such as vapour phase decomposition of zirconium-containing solutions (Sproson and Messing, 1984). Zr0 2

+

Y203

Mixing

I

ZrOCI2 0H2O + YCI3 Coprecipitation

I

C3H7OH

Melting

Filtration , Washing

Zr(OC3H7)4

Crushing , Milling

Drying

x

Hydrolysis

I

PSZ-Powder

I

Calcination 800 - 900 ° C

Z r 0 2 - Gel

Deagglomeration

Calcination

___

j

ZKX, / Y 2 0 3 TZP - Powder

Fig. 13.

TJ

Zr- Metal

Z r 0 2 powder production processes.

ZrOn

Processing and Properties of Advanced Structural Ceramics

77 %

0,05>jm i

(b) Fig. 14. Coprecipitated Zr0 2 /3 mol% Y 2 0 3 powder: (a) as precipitated; (b) after calcination at 850°C

78

L J. Cauckler

However, all the powders available on the market at present for the TZP ceramics are manufactured using precipitation techniques. These powders have almost uniform particle sizes of about 20-100 nm, and some are free from large hard agglomerates. Silicon

nitride

Silicon nitride powders are manufactured using a number of techniques, such as nitriding of silicon powder (Schwier, 1984), by chemical vapour deposition of silicon chloride and ammonia (Kijma, 1983), by the carbothermic reduction of silicon dioxide in a nitrogen-containing atmosphere (Zhan and Cannon, 1984) and by silicon diimide precipitation from silicon tetrachloride and subsequent calcination (Yamada and Kawahito, 1983). The most economical production of the powder is achieved by directly nitriding silicon. This technique, however, does not yield material of the highest purity. It has been found that the highest purity powders come from gas-phase reactions. As pure silicon nitride powders are very dificult to sinter and metal oxides are used as densification aids, some silicon dioxide impurities can be tolerated. These impurities, however, must be consistent in nature and quantity from batch to batch. An advantage of the silicon chloride and the silane route is that if the process manufacturing conditions are optimized, a very fine powder can be produced which requires no subsequent milling. The requirements for a high-quality silicon nitride powder suitable for sintering are fine particle size, constant low oxygen content, low free silicon and carbon impurities and high a-phase content. Wotting and Ziegler (1986) have examined the influence of the powder properties on pressureless sintering of a range of silicon nitride powders which were manufactured by the processes listed above. However, because of a number of complicating factors it was not possible to arrive at a general conclusion as to which was the best manufacturing process for the powder. Of much greater influence than the method used were the oxide and nitride additives needed as densifying additives. Silicon

carbide

Silicon carbide may be synthesized by many process routes (Gmelin, 1986), the most common of which is the Acheson process. This involves carbothermic reduction of silica and produces a-SiC material. Although normally the batch process involving high temperatures (>2000°C) is used, powders can also be produced carbothermically in a continuous process (Ibigawa Denko Co., 1976) at lower temperatures. The resulting product, however, is a j3-SiC

Processing and Properties of Advanced Structural Ceramics

79

powder. In both cases the product is milled to obtain submicron-sized particles. It is generally accepted that a-SiC powders show less tendency to excessive grain growth during sintering than /?-SiC powders. Table 4 summarizes the characteristics of the different methods. Other methods employ pyrolysis of polycarbosilane products or gas phase reactions of silicon compounds such as silicon chloride. Powder processing and

forming

Powder processing is very closely related to the forming process of the part itself. Components can be manufactured using dry pressing of powders, tape casting, extrusion, injection moulding, slip casting and pressure slip casting. Uncontrolled agglomerates lead in all methods to serious structural faults and these always result in unreliable components (Fig. 15). Starting from the spray-dried powder with hard agglomerates it is possible to see in the green body pressing defects which take the form of inhomogeneous pores; these lead to clearly discernible defects in the sintered component. The object of every powder preparation is therefore to produce synthetically a submicron powder that is ideal for the subsequent forming method, e.g. for dry pressing a controlled soft agglomerated form, free of coarse hard agglomerates, with homogeneously distributed organic additives. For all methods it is desirable to produce a homogeneous green body with high green density. These basic requirements have been the subject of much work in recent years. Pampuch and Haberko (1983) have provided Table 4.

SiC manufacturing routes.

Manufacturing route

Properties BET (m 2 g- 1 )

o2 (%)

Advantages

Disadvantages

Carbothermic reduction + Comminution Si0 2 + 3C -► SiC + 2CO (Acheson)

10-15

0.3-0.8

Low cost, high production rate, a-phase

Low purity, milling, irregular shape

Pyrolysis CH3SÌCI3 -► SiC + 3HCI

22

0.05

No milling, high purity, globular shape

Expensive, low production rate, ß-phase

Gas phase reaction SiCl4 + CH4 -► SiC + 4HCI

9

0.1-0.3

No milling, globular shape

Expensive, low production rate, impurities (CI,Si), /?-phase

L ). Cauckler

80

330 im i

<

(a)

100 ym ι

' -'■■I

(b)

100 \m t

t

(c)

Fig. 15. Developments of defects during dry pressing of spray-dried powder; (a) spray dried; (b) cold isostatic pressed (100 MPa); (c) sintered.

information about the risk of uncontrolled agglomerates on the sintering structure of fine powders. Especially significant is the danger of agglomerate formation with fine ceramic powders, such as zirconium oxide for TZP applications. Roosen and Hausner (1987) have shown how this agglomeration can be avoided by the spray freezing and spray drying of powders in the last stages of powder manufacture. Spray roasting is also employed. The chemicalmechanical steric fixing of individual particles which are manufactured using the sol-gel process has also been suggested (Johnson and Gallagher, 1978). Lange (1984) carried out a systematic examination of the causes of defects in ceramic structures. He also attempted, using a step-by-step procedure, to achieve their elimination. As an example an alumina-zirconia composite ceramic (Lange, 1986) was first dry pressed and subsequently achieved a strength of 600 MPa. The strength, however, was less a property of the material itself, but rather produced by the defects caused by agglomerates in the starting powders. In a subsequent step the same material was produced without agglomerates in the powder, and a strength of 1000 MPa was achieved. An analysis of the structure showed that the strength could be further increased by eliminating the fine pores from the structure. By further improving the distribution of organic additives and controlling their burnout, additional strength improvements were possible to 1400 MPa. Finally by hot isostatic pressing after sintering, strength could be pushed up to 2000 M Pa. This example clearly demonstrates the potential resulting from

Processing and Properties of Advanced Structural Ceramics

81

the improvement of the conventional powder route, enabling fabrication of high-strength reliable ceramics. Consolidation by sintering

The consolidation of porous powder compacts is achieved by means of solid phase or liquid phase sintering. During the solid phase sintering atoms diffuse away from the grain boundaries at the contact points of powder particles to their solid-gas interface. By these means the centres of the individual particles approach each other, and a shrinking process in the powder compact takes place. Concurrently, surface diffusion also takes place resulting in coarsening of the grains without densification. The surface energy of the powder, which is the driving force for sintering, is then used up in grain coarsening and does not give rise to any consolidation effect. During liquid phase sintering a part of the powder compact becomes liquid and allows rapid diffusion of atoms away from the grain boundaries with consequent rapid consolidation. The sintering process is generally accelerated by use offineuniform powder particles with high packing density and the use of high pressure and temperature. For the control of consolidation and the suppression of grain growth many techniques are available, most of them relying on dopants such as MgO in A1203. Knowledge on this subject is well advanced and has been reviewed recently by Harmer et al (1986). Recently new techniques have been introduced with the objective of controlling the competition between consolidation and grain growth during sintering in such a way as to give consolidation an advantage. This is achieved by rapid heating of the material as shown by Brook (1982) and Harmer and Brook (1981). In conventional heat treatment, however, this is limited by the relatively slow heat transfer and by the thermal shock sensitivity of the green body. By rapid heating with microwaves in the centimetre range it is possible to sinter very satisfactorily aluminium oxide components. The experiments of Janney and Kimrey (1988) have given very encouraging results (Fig. 16). Alumina from Sumitomo (AKP50) was heated both by conventional means and by use of microwaves (286 GHz). Use of the microwave technique made it possible to achieve 97% of theoretical density at temperatures as low as 1100°C. In contrast, the conventionally heated samples showed only 69% of theoretical density. Tian et al. (1987) carried out similar experiments and also obtained very uniform microstructures in compacts with Linde CR30 alumina powder. Test pieces with an average of 0.8-0.9 μπι grain size were

L J. Gauckler

82 100

Temperature (°C) Fig. 16. Microwave (MW) versus conventional sintering of alumina (Brew) (Janney and Kimrey, 1988)

obtained in contrast to the normal 3-7 μπι range in rapidly but conventionally heated material. Composite materials made of alumina-titanium carbide have also been successfully consolidated using this technique (Tian et ai, 1987). The advantage of using microwave heating is obvious. The thermal energy can be applied directly into the area required without being limited by thermal conductivity and size of the compact. In addition, the application of energy by microwave heating can be matched to the product geometry. There are, however, some disadvantages which must be taken into consider­ ation. Ceramic materials have differing electromagnetic properties depending not only on composition, but varying also as a function of temperature. Thus the necessary hardware for the sintering process must be adjusted to suit the customer's requirements. A number of specific technical problems remain to be solved, such as insulating materials and temperature control inside the compact. It is suggested that microwave heating in combination with other conventional heat sources offers possibilities for further develop­ ment (Tinga, 1986). Enhanced gas pressure during or after sintering has been in use for some time to improve the density of superalloys and ceramic materials (Hunold, 1983). It is well known in the case of ceramics that hot isostatic pressing

83

Processing and Properties of Advanced Structural Ceramics

(HIP) using pressures in the range of 100-200 MPa at temperatures between 900 and 1900°C achieves structures with virtually no residual porosity and without any detectable massive grain growth. In addition, there are further advantages in the manufacture of ceramics by use of HIP techniques. Sintering aids for example can be drastically reduced, and a fine-grained microstructure can be obtained. Materials, which have a tendency to decompose, can be compressed at lower temperatures. The main advantage, however, results from the reduction in the spread of mechanical properties which is achieved by elimination of the larger defects in the internal structure. It has been very clearly demonstrated by Neil et al (1988) that the increase in reliability in various process steps can be achieved by use of enhanced pressure. In this work, the starting material was silicon nitride containing 6wt% yttrium oxide and 2wt% alumina. The largest defects in the isostatically pressed green body could be shown by non-destructive testing to be density inhomogeneities in the green compact as well as inclusions of silicon particles. The sintered material had a strength of 697 MPa and a Weibull modulus of m = 9. After eliminating these defects by improved processing and by using injection moulding and sintering, it was possible to obtain materials with strengths of 800-900 MPa in combination with a Weibull modulus of m = 14. Post-sintering and use of H^P techniques on this material have pushed strengths up to almost 1000 MPa with a Weibull modulus of m = 20 (see Fig. 17).

500

700

900

1100

Modulus of Rupture (MPa) Fig. 17.

Increasing reliability of Si3N4 ceramics by improved pressing (Neil et ai, 1988).

84

L j. Gauckler

The latter methods clearly demonstrated that much more reliable structural ceramics are possible by improved processing avoiding broad defect distri­ butions with large defects. All these methods, however, need additional or expensive processing steps which increase the costs of these materials. Therefore more work is needed towards highly reliable as well as low-cost structural ceramics.

Alternative manufacturing procedures

As an extension to the classic powder metallurgical processes a number of alternative manufacturing methods are explored today. Some of them aim at the production of perfect monolithic ceramic structures. A good example of this is the sol-gel process (Hench et al, 1988). Another recently explored process, the melt-oxidation process, has as its objective the near-net-shape production of ceramic composites with increased fracture toughness. (Lanxide process, Newkirk et al, 1987). Other techniques, which are used for the production of composites, are slurry infiltration of fibre prepregs and preforms with subsequent densification (Brennan and Prewo, 1982). This process is mainly used for carbon and silicon carbide glass fibre composites. It can be combined with the sol-gel impregnation process of fabrics. Further examples are provided by the pyrolysis of polymers in which SiC/SiC composites are manufactured (Fitzer and Gadow, 1986). Yet another tech­ nique is the impregnation of fibre materials via the gas phase (Naslain and Langlais, 1986). The sol-gel

technique

(see also Chapter 3)

The sol-gel technique appears to be a most elegant method of overcoming the disadvantages of the powder metallurgical manufacturing processes. This well-known technique, which has already been applied for nuclear fuel spheres, functions from both colloidal solutions and metal-organic combi­ nations, such as alkoxysilanes (Johnson, 1985; Colomban, 1985). In the case of metal alkoxides, they can be reacted with water to produce a finely distributed sol (Hench and Lee, 1987). This latter technique was developed mainly for the production of ceramics based on silicon dioxide. The sol is cast into a mould, gelled, dried, and the extremely fine pored compact is sintered at temperatures 200-300°C lower than conventional powder compacts. The advantage of the sol-gel process lies in the fact that homogeneous mixtures on the atomistic level are obtained by mixing the alkoxide, and subsequently a very fine-grained homogeneous green body is achieved which

Processing and Properties of Advanced Structural Ceramics

85

can be sintered at temperatures as low as 1000°C (Klein et a/., 1984). The forming of complex shapes, including the spinning of fibres, is also possible at low temperatures. The major disadvantages are the cost of the metal alkoxides, the high shrinkage during dehydration of the gel as well as the long manufacturing time. It seems that this particular technique is best for films, fibres, thin layers and small thin-walled components as well as for optical applications (Hench et ai, 1988). It offers, however, so many advantages that development of larger components is also in progress as demonstrated in Fig. 18 (Hench et a/., 1988).

Fig. 18.

Silica optical components made using the sol-gel process (Hench et ai, 1988).

Direct melt

oxidation

Another very interesting manufacturing method is the Lanxide process (Newkirk et al, 1986). The process involves the formation of ceramic-metal composites by the oxidation of molten metal such as aluminium. At temperatures between 1400 and 1650°C, an oxide product on the melt resulting from the liquid-gas reaction forms a coherent ceramic with pore channels enabling further liquid metal transport from the melt through the ceramic product for further oxidation. This allows growth of a composite of uniform microstructure with metal contents of 5-30 vol% as long as molten metal and oxidant are available. Particularly interesting is the formation of matrix composites via this process. The matrix can be formed by placing an inert filler material (e.g. fibres, whiskers, particles or platelets) as a formed powder compact adjacent to the parent metal surface in the path of the outward oxidation growth process (Fig. 19). During oxidation the reaction

L ). Cauckler

86 Vapor Phase Oxidant

Vapor Phase Oxidant

Filler With Reaction Product Matrix

ty&zzzzzy

V&zzzz$

Refractory Container

Refractory Container

a)

b)

c o o

Af-3*toSi-i

CO

c "<5

(5

O) Φ

o

"o

Φ a (/)

1350

1400

1450

1500

1550

1600

1650

Setpoint Temperature (K)

c) Fig. 19. Direct melt oxidation: a, without filler; b, with filler, c, Effect of Mg concentration on specific weight gains following 24 h at different setting-point temperatures of Al and Al-Si alloys (Newkirk et ai, 1986, 1987).

Processing and Properties of Advanced Structural Ceramics

87

product grows through the filler. Normally, the filler material (the shaped component) is not destroyed during the oxidation, and therefore near net shaping is possible with no or little machining necessary on the final component. A variety of ceramic-metal systems has been demonstrated including AlA1 2 0 3 , A1-A1N, Ti-TiN, Zr-ZrN and others. In addition a range of fillers are in use, among them A1 2 0 3 or SiC fibres and woven cloths. The cost factors for such composites seem to be very attractive. In addition, there is no grain boundary material involved during fabrication of these materials (e.g. in A1 2 0 3 ), thus it is expected that they will retain their strength at elevated temperatures. This new technology can be applied to making large ceramic bodies because densification shrinkage is eliminated (Newkirk et al, 1987). Impregnation

of fibre

structures

Fibre-reinforced ceramic composites are manufactured using fibre prepregs impregnated by slurries. After drying the structures are sintered or hot pressed. This technique is suitable for lithium alumina silicate (LAS) systems with, for example, carbon or silicon carbide fibres. The technique is less suited for systems with high refractory matrices, since during sintering at high temperatures undesirable fibre matrix reactions can easily occur (Cornie et a/., 1986). This disadvantage can be overcome in part by impregnating the fibre network by means of the sol-gel method followed by low-temperature sintering. This gives a better homogeneity of the matrix, easier infiltration and allows lower sintering temperatures to be used. These in turn do less damage to the fibres themselves. One disadvantage, however, is the relatively large shrinkage which occurs during densification. This can be somewhat reduced if the sol is combined with a powder slurry. Another method of direct forming of composites involves the pyrolysis of organosilicon compounds. Liquid organosilicon monomers or polymers are used to impregnate for example SiC fibre structures. The first step involves production of a porous fibre skeleton impregnated with a small amount of binding phase which is then evacuated in an autoclave and impregnated with liquid organosilicon compounds at 500 to 600°C under decreasing N 2 gas pressure up to 40 M Pa. At increasing temperatures the oligomer silanes are transformed to polycarbosilanes and polymerized. After repeating this step several times, the material is decomposed in an autoclave at 1000°C to SiC. It has been shown that thermal decomposition of polymeric organosilicons can also be used to form powders (Mazdiyasni, 1982), fibres (Yajima, 1976; Tanaka, 1980) as well as monolithic ceramics of silicon carbide and silicon

88

L. J. Gauckler

nitride (Mazdiyasni et al, 1978). Preparation and properties of monolithic and composite ceramics pro­ duced by polymer pyrolysis were reviewed recently by Rice (1983). In his survey he covered various precursors with cage-, phenyl- and unsaturatedring structures. Decomposition yields as high as 60% were achieved. Small bulk pieces of ceramics were produced with fine microstructures. Strength values, however, were medium to low. The method today seems to be suitable for fibre production, but it is less developed for bulk materials. Another promising method, again especially suited forfibre-reinforcedSiC materials, is the in situ chemical reaction technique, such as chemical vapour deposition and chemical vapour infiltration. This involves using a mat or a three-dimensional fibre structure of for example SiC fibres, which is vapour impregnated by CH3SiCl3. This is thermally decomposed to SiC at 12001400°C (Naslain, 1981; Naslain and Langlais, 1986). SiC/SiC composites thus show 10% open porosity, but they have enhanced toughness with high fracture energy values of up to 10 000 Jm~2,'corresponding to toughness values in the range of 25 M Pam* (Lamicq et al, 1986).

IMPROVED PROPERTIES BY MICROSTRUCTURAL DESIGN

Toughness of ceramics can be improved by microstructural design methods. The most important of these are: dispersion of a second phase (Wei and Becher, 1984; Terao, 1988), transformation toughening (Garvie et al, 1975; Claussen and Riihle, 1981; Evans and Cannon, 1986), as well as reinforcement of a matrix by fibres and whiskers (Rice, 1983; Cornie et al, 1986; Vaughn et al, 1987). In addition, high toughness can be achieved in systems with ceramic grain networks embedded in glassy matrices by proper selection of thermal expansion of the glass. In the case of A1203 grains bonded by a 15wt% grain boundary glass, the higher coefficient of thermal expansion of A1203 creates microscale residual thermal compressive stresses in the glassy phase and corresponding tensile stresses in the A1203 skeleton. This leads to extensive cleavage fracture of the A1203 grains resulting in toughness as high as 8.5MPam* compared to 3.5MPam* for comparable compositions where the crack propagates preferably through the grain boundaries (Travitzky et al, 1985a,b) if the glass is partially crystallized. Compressive stresses can also be incorporated in the surface of a component by grinding (Gupta et al, 1978; Claussen and Riihle, 1981; Claussen, 1983), by ion implantation or by glazing the surface. Some reviews published recently show the high level of scientific and technical interest in this field (Chawla, 1983; Fisher, 1984).

Processing and Properties of Advanced Structural Ceramics

89

Crack deflection by a dispersed second phase

Two-phase ceramics are in general less brittle than single-phase materials. An increase of 50% in the fracture toughness Klc can be achieved in this way. The reason for this behaviour is crack deflection during fracture. The crack will be deflected parallel to its axis by the second phase particles, or the crack plane is rotated. To induce rotation of the crack plane and thus increase toughness, elongated particles are more effective than globular particles. Faber and Evans (1983) have established a model for the effect of crack deflection around second phase particles on the fracture toughness of brittle materials. Some examples of ceramics, in which toughness is increased by the introduction of a second phase, are shown in Fig. 20. This composite approach has been developed for many systems such as SiC-TiB2 (McMurtry

Fig. 20. Toughening by crack deflection in ceramic/ceramic composites (McMurtry ei ai, 1987; Rice, 1983; Ruf and Evans, 1983; Terao, 1988). For abbreviations see text.

90

L J. Cauckler

et al.9 1987), SiC-TiC (Wei and Becher, 1984), A1 2 0 3 -BN (Rice, 1983), Z n O Z r 0 2 (Ruf and Evans, 1983) and S i 3 N 4 - Z r 0 2 (Terao, 1988). The second phase in the microstructure also can prevent uncontrolled grain growth of the matrix phase during sintering. However, it is important to have particles with a subcriticai diameter. Particles that are too large introduce stresses which are too high during cooling of the compact, leading to the development of microcracks. This can also contribute to a toughness increase. Microcrack toughening can be introduced into a microstructure by two processes; (1) localized residual tension because of the mismatch of thermal expansion between the grains and the matrix, or (2) by a phase transformation of the particles (Claussen et al, 1982). In the first case, a system with second phase particles in a rigid matrix with a larger thermal contraction coefficient than the host tends to produce microcracks around the particles (Evans and Faber, 1981). In the second case, the second phase undergoes a martensitic phase transformation during cooling from high temperature. The associated volume expansion creates tangential stresses which can cause microcracking of the matrix around the particles if the particles are larger than about 1 μηι. The strain energy of the crack is lowered because the microcrack interacts with the main crack (Claussen et ai, 1982; Butler, 1985). It is important for both kinds of microcracking, which in most cases cannot be separated, to have a uniform spacing, a narrow size distribution and an angular morphology of the second phase particles (Evans and Faber, 1981). It is observed that microcracking in the A l 2 0 3 - Z r 0 2 system decreases hardness, but the critical stress intensity factors of 9 lOMPam* were much higher than for A1 2 0 3 alone (Claussen, 1976; Green, 1982).

Transformation toughening

Increased toughness of ceramics by stress-induced phase transformation was first discovered by Garvie et al (1975) in the system Z r 0 2 and applied by Claussen (1976) to the system A l 2 0 3 - Z r 0 2 . This improvement in toughness relies on the stress-induced martensitic phase transformation of tetragonal Z r 0 2 or Hf0 2 to the monoclinic form with a volume increase of 3 5vol%. It is important for transformation toughening that the stable hightemperature tetragonal modification is retained as a metastable state down to room temperature. The phase transformation then can be triggered by stresses like those in front of a crack tip. On further loading the crack tip has to advance into the transformed region offering resistance to crack propagation, resulting in increased toughness. The increase of the fracture energy by this crack shielding depends on the amount of transformable

91

Processing and Properties of Advanced Structural Ceramics

particles (grains) in the microstructure and the width of the transformation zone around the crack (Evans and Cannon, 1986). The science of transfor­ mation toughening has developed well in the last few years (Evans et al, 1981) and has been the subject of several conferences. The tetragonal phase of Zr0 2 and Hf0 2 can be stabilized to lower temperatures by solid solution formation together with metal oxides such as Y 2 0 3 , rare-earth metal oxides, MgO, CaO and Ti0 2 . In addition, the transformation temperature is lowered further by a reduction of grain size. VeryfineZ r 0 2 grains (0.1-1 μπι) doped with 3 mol% Y 2 0 3 remain metastable in the tetragonal form far below room temperature. Three different types of Zr0 2 transformation-toughened materials are known (see Fig. 21). (1) Single phase Zr0 2 of tetragonal zirconia polycrystals (TZP) shows rather high toughness (6 to 10 M Pa m*) and very high strength ( > 2500 M Pa for material with A1 2 0 3 additions obtained by hot isostatic pressing, Watanabe, 1983; Tsukuma, 1983). The grain size is 0.3-1 μπι. Both toughness and strength depend on the Y 2 0 3 content and the grain size. (2) A second possibility of toughening is by producing fine coherent tetragonal precipitates in a cubic zirconia matrix. The microstructure of this material (partially stabilized zirconia, PSZ) shows the typical large cubic grains with fine tetragonal precipitates. These are formed during annealing the material with composition in the two-phase field, cubic and tetragonal, below manufacturing temperature. PSZ materials have typical strengths of 1000 MPa and a Klc of 8-10 MPam* (Heuer and Hobbs, 1981). (3) In the third case, the small zirconium oxide particles of tetragonal

(a)

(b)

(c)

Fig. 21. Microstructures of transformation-toughened ceramics: a, tetragonal zirconia polycrystals (TZP); b, partially stabilized zirconia (PSZ); c, zirconia-toughened alumina (ZTA).

92

L J. Gauck/er

modification are dispersed in a ceramic matrix, e.g. A1203 (TZA) (Claussen, 1976), mullite or other chemically compatible matrices. In situ straining experiments were performed on A1203 containing 15 vol% Zr0 2 , in a transmission electron microscope by Riihle et al (1983). They observed the transformation of tetragonal Zr0 2 particles to monoclinic symmetry close to the crack tip during straining. In several of these directly observed crack propagation experiments Riihle et al. (1983) clearly demonstrated the presence of the transformation zone ahead of the crack tip. The width of the transformation zone was evaluated to be several microns, as shown in Fig. 22 by the dark (monoclinic) particles around the crack in contrast to the light grey particles which make up dispersed untwinned tetragonal zirconia phase. By embedding Zr0 2 particles in matrices, a whole range of ceramics have been improved (see Fig. 23). The chemical compatibility between matrix and dispersed Zr0 2 is one main requirement for high toughening results (Claussen, 1983). The critical particle size of Zr0 2 varies with different matrices and is between 0.5 and 1.2 ^m. Application of transformation toughening is limited to temperatures below 800°C. Some of these materials also show fatigue damage under cyclic loading. Claussen (1985) pointed out the possibilities of extending transformation toughening to higher temperatures.

Fig. 22. Transmission electron micrograph of 1 5 v o l % Z r 0 2 (tetragonal and monoclinic) toughened alumina w i t h a crack. The dark particles in the process zone around the crack are twinned monoclinic Z r 0 2 . The light grey particles are tetragonal Z r 0 2 . (Courtesy of M. Rühle.)

Processing and Properties

of Advanced

Structural

93

Ceramics

16

14

15

Vol % Z r 0 2

12

i,.

22

03 0.

25 8 15

C/)

CD C H

6-I

3

4

o

r-

20

2

Ü

(?) I

ÛL

c N

Z

CO

(?) I

z

CO

(75

OCM

< I

o

CM

<

I I Û. Û. 0) (/) X Q. X Z r 0 2 transformation-toughened ceramics. For abbreviations see text. X

X

Fig. 23.

o I

Fibre and whisker composites

Fibre reinforcement of ceramics holds out the promise of making them less susceptible to damage and preventing brittle failure. Fibre-reinforced ceramics are basically different from metals. In the case of metals, high-strength ceramic fibres with a high elastic modulus are embedded in the matrix to stiffen the metal, i.e. to increase Young's modulus, often at the cost of toughness. In ceramic materials, fibres are incorporated into a matrix to increase toughness, sometimes yielding to a lower bend strength. The increased fracture toughness is due to crack shielding in the process zone, enlargement of the surface of the fracture, crack deflection, crack branching and crack bridging and a debonding of the fibre-matrix interface with subsequent fibre pull-out (Marshall and Evans, 1985). Thus, the stress-strain curve offibre-reinforcedceramics shows a pseudoplastic behaviour after matrix cracking. Furthermore the fracture behaviour of these ceramics can no longer be characterized by a single value Kc, but rather by the work of fracture. The development of fibre- or whisker-reinforced ceramics is not new, but

94

L. j. Gauckler

it has regained attention in recent years. Great progress has been achieved in the development of lithium aluminium silicates (LAS) reinforced with fibres (Brennan and Prewo, 1982). One reason for the renewed interest is the availability of better and more oxidation-resistant SiC fibres. Today yarn is made out of SiC, with and without a carbon core, in addition A1 2 0 3 , A1 2 0 3 borosilicate yarn as well as α-Α1 2 0 3 and A l 2 0 3 - S i 0 2 fibres and SiC whiskers are also available. See Table 5. In addition to improved toughness, high strength is also desirable. The conditions for improving strength of fibre-reinforced ceramics relative to normal ceramics is that the fibre spacing is smaller than the flaw controlling the fracture of the matrix. It is further important that high (> 50vol%) volume fractions of fibres are distributed uniformly in the matrix. Both conditions are not easy to fulfil. Therefore it is easier to increase the toughness of the fibre-reinforced ceramic than to increase its strength. The increase in toughness depends mainly on the interfacial bond between the fibre and matrix. If the bonding between the fibre and the matrix is too weak, the materials will not transmit compressive and shear stresses. On the other hand, the material shows brittle fracture when the bonding is too strong. A deflection of the crack at the fibre-matrix interface is necessary to avoid brittle fracture. See Fig. 24 for the toughness and strength of fibrereinforced ceramics and glass. Another condition for the successful combination of a ceramic and a fibre is that the thermal expansions of fibre and matrix are matched; otherwise spontaneous cracking can occur during manufacture or in service. Today fibres are incorporated in a matrix in an oriented way, as a tissue, or undirected as chopped fibres. Manufacturing processes of fibre-reinforced ceramics include pressing of powder fibre mixtures, the sol-gel process, infiltration by a liquid phase, gas infiltration and pyrolysis of organic ceramic precursors such as polysilanes. The different manufacturing processes have recently been reviewed by Cornie et al (1986). A large number of fibre and matrix systems has been developed on the laboratory scale. Only a few examples will be mentioned here (Lamicq et al, 1986). Whisker-reinforced materials in the system Al 2 0 3 -SiC were prepared by Becher and Wei (1984), having a Kic of 12MPam*. Improved toughness is obtained with very fine, large whiskers. Such materials keep their high fracture energy even at high temperatures. Creep at high temperatures is also reduced by this whisker reinforcing. The whiskers have a diameter of 0.5-3 μηι and an aspect ratio of 200. Hot pressed Si 3 N 4 with SiC whiskers were investigated by Shalek et al (1986). For two-dimensional laminates of SiC/SiC composites Lamicq et al (1986) showed that toughness can be increased up to 25 M Pa m*, corresponding to a fracture energy of

Properties of fibres and SiC whiskers.

SiC on carbon core (SCS 6, Avco) SiC-Nicalon (Nippon Carbon) FP-AI 2 0 3 (El DuPont de Nemours) AI203-borosilicate yarn (3Μ, Nextel 312) A l 2 0 3 - Si0 2 yarn (Sumitomo) VLS-SiC whisker (Los Alamos Nat. Lab.)

Table 5.

390 150 230 580

1.5 1.7 8.4

1.8-2.6

180--210

406

Young's modulus (GPa)

2.5-3.5

3.9

Tensile strength (GPa)

6

9-17

20 10

10-20

140

Diameter (/im)

3.0 2.6 3.9 2.5 3.2 3.3

Density (gern" 3 )

96

L J. Gauckler

1500-1 TO

5 O)

1000J

o 3

'S

D

500 H

LL

Ü

δ

Ü

δ Ü

IL· 1

Ü Ü c/5

c/5

<7>

Ü

z

d


< a.

5

X

Q.

+

55

+

CO

X

+

CI

+ en I

X Fig. 24. Toughness and strength of fibre-reinforced ceramics and glass. For abbreviations see text.

10000 J m ~ 2 at a bend strength of 380MPa. This material is manufactured by gas phase infiltration of Nicalon (from Nippon Carbon) SiC-fibre tissues. An improvement of lithium aluminium silicate glass (LAS) by SiC-coated carbon fibres, produced by slurry impregnation and hot pressing, is described by Brennan and Prewo (1982). Hot pressed Si 3 N 4 reinforced by SiC fibres with bend strengths of 500-600 MPa and a toughness up to 12.5 MPam* is reported by Shalek et al (1986). Combinations of fibre reinforcement and other toughening mechanisms, e.g. transformation toughening in the case of TZP and mullite, are also possible. A summary of today's efforts is given by Menges and Ziegler (1988). Not all the properties desired can be attained by these material combi­ nations. This is because of insufficient knowledge of the high-temperature behaviour. Nicalon fibre composites, for example, show inadequate oxidation resistance if the matrix is cracked (Mah et a/., 1987). Even without any

Processing and Properties of Advanced Structural Ceramics

97

cracks, a so-called pipeline oxidation along the interface between matrix and fibre can occur. Another drawback is the high cost of manufacturing. For SiC fibrereinforced SÌ3N4, gas pressure sintering or HIP in N 2 atmosphere is needed. Even for large-scale manufacturing of fibre-reinforced ceramics with continuous fibres, the production technologies are complex and expensive for applications outside the high-technology area.

Grain boundary engineering

Si 3 N 4 is one of the most promising advanced structural ceramics with already many applications in the medium- and high-temperature region. These materials are sintered with the aid of an oxide forming a refractory silicate grain boundary glass. Dense Si 3 N 4 has a microstructure with elongated grains which are responsible for a high toughness of up to 6MPam*. This material also shows good high-temperature strength (600-800 M Pa) up to 1250°C. At higher temperatures the lifetime is limited by creep deformation for medium loads, caused by the softening of the silicate, while in the higher loading range the lifetime is limited by slow crack growth accompanied by the formation of pores in the grain boundary triple points. Both processes depend on the glass phase, which is always present in the grain boundaries of these materials (Hofmann and Gauckler, 1974; Rühle et a/., 1977). For the improvement of high-temperature properties of dense silicon nitride, there are three strategies. One possibility is to replace the glass phase necessary for sintering by a more refractory one. Gazza (1975) showed that Si 3 N 4 materials with Y 2 0 3 containing glass had higher strength at high temperature compared to materials with MgO containing glasses at the grain boundaries. Another approach to improving the high-temperature mechanical proper­ ties of silicon nitride is by forming solid solutions of ß-Si 3 N 4 together with an oxide densification aid. This is possible if the atoms of the densification oxide are incorporated in the Si 3 N 4 lattice at high temperatures after densification has taken place. It was discovered first by Oyama and Kamigaito (1971) and then by Jack and Wilson (1972) that Si 3 N 4 forms solid solution with A1 2 0 3 . However, it was not possible to reproduce single-phase Si—Al— O-N materials unless the phase equilibria in the system Si-Al-O-N were well established (Gauckler et a/., 1975; Huseby et a/., 1975). Based on these results, a transient liquid-phase sintering process could be developed (Gauckler et al, 1978) using A1 2 0 3 and A1N as well as S i 0 2 as liquid-forming constituents. The nitrogen-containing glass phase enhances rapid densification during heating up. The elements of the glass are

Fig. 25. Si3N4 densified with Al 2 0 3 and Y 2 0 3 additives, a, Dark-field image formed with electrons diffusely scattered by the glass phase (light), b, Dark-field image formed with YAG reflection showing crystallized YAG grain boundary phase (light). (Courtesy of D. Szabo, MPI, Stuttgart.)

99

Processing and Properties of Advanced Structural Ceramics

subsequently absorbed by the j?-Si3N4 grains forming a solid solution. Materials with very little glass in the grain boundaries can be produced this way. Lumby et al (1978) gave the evidence for drastically improved hightemperature creep resistance of these solid solutions compared to materials with substantial amounts of grain boundary glass. Transmission electron microscopy observations of these creep-resistant materials showed that after the formation of solid solutions the grain boundaries still contained glass phases of a thickness of some nanometres. Clarke (1987) reasoned that the glass film in grain boundaries of Si3N4 might have an equilibrium thickness of 5-6 nm. It was further shown that the high manufacturing temperatures of the solid solutions in the pure SiAlON system caused the Si3N4 grains to become more globular and no longer elongated, which was detrimental to the toughness of Si3N4 (Gauckler and Petzow, 1977; Gauckler et al, 1977). Therefore, a third way is chosen today to improve creep resistance of sintered silicon nitride. To sinter silicon nitride, an yttria- and alumina-containing glass is used. In a post-sintering heat treatment the glass then can be crystallized. This procedure was proposed by Tien (personal communication, 1978) for silicon nitride with yttrium-aluminium-garnet (YAG) and cordierite as sintering aid (see Fig. 25). The controlled crystallization of grain boundary glass during post-sintering was carried out in these systems by Bonneil et al (1987). Chen and Tien (1987) reported improved high-temperature creep resistance (see Fig. 26). 500 010 7G SIALON

400

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100

L J. Gauckler

The microstructure of these materials lying in the SiAlON-Y-garnet systems consists of large areas of garnet crystals with approximately identical orientations in the grain boundaries after crystallization of the glass. These materials with crystallized grain boundary phase show less deformation at high temperatures.

Summary

The development of advanced structural ceramics is determined by the correlations between production technology and microstructure as well as between microstructure and properties. It appears that for particular functions only a special combination of production process and material is suitable. By coordinating the production processes with the application functions of these materials, great advances can be expected. This will open up an increasing number of applications for advanced structural ceramics in mechanical applications. Whereas ten to fifteen years ago most ceramics were almost exclusively used under compressive load, new developments in strength and reliability now allow these materials to be used also under tension. The recent development in the mechanical properties of these materials, shown in Fig. 2, is impressive. The strength values of 1-2 GPa achieved today are not the limit. One of the main aims of further development will be to translate new knowledge into producing special microstructures of highly reliable ceramics in a broader range of materials, on a daily basis. Together with the improvements in classical powder metallurgy pro­ duction, also very promising is the search for low-cost alternative manufactur­ ing processes. From the formulation of new composites, further improvements can be expected. A good example of microstructural design is the consequent development of transformation-toughened ceramics. Whisker- and fibrereinforced ceramics have just recently shown their great potential. The increasing awareness of the unique properties of this class of materials and their wider use will also lead to more precise and suitable design methods. Especially in the case of new concepts the consistent application of design criteria will pay off through increased chances of success. This will also widen the data base of properties which is far from complete. There is no doubt that advanced structural ceramics will be involved in the realization of new heat engine and energy concepts. Whether they will be used in large quantities in automotive heat engines and gas turbines and whether they will replace metals in these applications is still an open question today.

Processing and Properties of Advanced Structural Ceramics

101

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