Steel Metal

Introduction

Carbon-graphites offer the design engineer a unique family of mechanical materials. Manufactured entirely from carbon and including high temperature carbonaceous bonding, these materials combine the strength, hardness and wear resistance of carbon with the corrosion resistance and self lubricating properties of graphite. The precisely controlled inherent porosity of carbon-graphites can be filled with a variety of impregnants to enhance chemical, mechanical and tribological properties.
Types of Carbon

The terms ‘carbon’ and ‘graphite’ are often used interchangeably. This is unfortunate since each form of the element carbon offers specific properties that can be used to benefit different types of applications.
Amorphous Carbon

Amorphous carbon is a very hard, strong compound. The crystals exhibit a turbostratic disorder which makes the material extremely resistant to wear. The strength and wear resistance properties of this material make it of interest in some applications. However, these strengths can also be a weakness -carbon generates high friction when rubbed against another surface.
Graphite

Graphite, on the other hand, is softer and relatively weak because of the crystalline order and closer spacing between the monoplanes and stacks. A graphite structure can be compared to a deck of cards with individual layers able to easily slide off the deck. This phenomenon gives the material a self lubricating ability which is matched by no other material. External lubricants are simply not necessary.
Carbon-Graphites

It is possible to combine amorphous carbon and graphite to take full advantage of the strengths and weaknesses of each of these two types of carbon, table 1. The proper mixture of the two materials is strong and hard and has low friction. At the same time, this composite has excellent corrosion resistance and is capable of operating at temperatures in excess of 315°C for extended periods of time, depending on the specific grade. The ability to create materials that have these properties is the basis of the manufactured mechanical carbon materials that perform well in difficult tribological situations such as pumps.

Table 1. Properties of typical carbon and graphite materials.

 
    

Carbon 100%
    

Carbon-graphite (70%-30%)
    

Carbon-graphite (30%-70%)
    

Graphite 100%

Apparent density (g.cm-3)
    

1.70
    

1.72
    

1.75
    

1.80

Hardness
    

100
    

85
    

65
    

40

Compressive strength (MNm-2)
    

300
    

208
    

145
    

55

Flexural strength (MNm-2)
    

62
    

62
    

52
    

28

Modulus of elasticity (GNm-2)
    

21
    

17
    

14
    

10

Thermal conductivity (W/m.°C)
    

5
    

9
    

12
    

85

Temp limit in air (°C)
    

315
    

315
    

315
    

455
Processing Carbon-Graphites

Carbon-graphites are created by combining the two forms of carbon with coal tar pitch. The coal tar pitch acts as a temporary binder that holds the two structures together during the compression moulding process in which near net shapes are formed. Following the moulding operation, the parts are sintered at temperatures high enough to carbonise the coal tar pitch. The result is a structure that is completely carbon bound and contains both carbon and graphite. This structure is extremely strong in compression and will not creep under load. The carbonisation of the temporary binder leaves holes in the structure - on a micro scale the sintered body is a black sponge.

The formation of holes during the processing of a carbon-graphite composite has various advantages. The holes can be filled with resins, metals, carbon, or inorganic salts, depending on the planned use of the material, table 2. These fillers serve to improve the strength, thermal conductivity and tribological characteristics of the material. Additionally, carbon-graphite can be sintered to an even higher temperature to convert the entire structure to graphite to provide especially good performance in very high temperature, high speed applications.

Impregnation
    

Application

Thermoset resin
    

General duty to 260°C in water, coolants, fuels, oils. Light chemical solutions, food and drug.

Antimony
    

Hot water, steam, light hydrocarbons.

Copper or silver
    

High pressure service to 20.7MNm-2.

Carbon
    

Highly corrosive environments.

Film formers (fluorides etc)
    

Extremely dry environments, vacuum or cryogenic.

Oxidation inhibitors (phosphates etc)
    

High temperature and/or high speed (to 538°C and 240ms-1).
Carbon-Graphites vs Traditional Lubricants

Carbon-graphites are used in a wide range of applications where traditional lubricating methods are not appropriate. For example, a typical oil lubricated bearing struggles at temperatures below -40°C because of the high viscosity of the oil. Above 200°C oils carbonise, making them abrasive and ineffective. Graphite bearings are capable of extended use at temperatures above 600°C.
Chemically Aggressive Environments

Chemically aggressive applications represent another application niche for carbon-graphite. For instance, sterilisation processes tend to leach the oil from the structure of an oil lubricated bearing. Also, solvents and radiation can break down lubricating oils, and low pressures can cause the oil to vaporise. Carbon-graphite materials are inherently stable and chemically resistant, making them ideally suitable for these types of application, figure 1.

Figure 1. Inside of a pump. The four vanes on the rotor are made of carbon-graphite, and are consequently self lubricating, temperature resistant and impermeable to gases and liquids.
High Load Applications

Lubricants are inappropriate in other applications for a variety of reasons. High loads can squeeze the lubricant from a surface. Without the hydrodynamic layer of lubricant, failure is imminent if the material cannot provide self lubrication. In some applications, such as those involving food handling, lubricants can contaminate the surrounding environment. Specific grades of carbon-graphites are approved for use in food handling applications.
Design Factors

Carbon-graphite bearings are used in both wet and dry operating conditions. Carbon-graphite allows the designer to specify the bearing close to the boundary lubricated condition without the risk of seizure. Permissible loads and running speeds depend on the allowable wear rate. Shaft materials and surface finishes are important factors in the wear rates of carbon-graphite materials. As a rule of thumb, the harder and more polished the surface, the lower the wear. Where possible, aluminium and bronze should be avoided for use as shaft materials.

Carbon-graphite materials are also widely used as rotating shaft and face seal materials, figure 2. They perform well when running against metal and ceramic counterfaces. Seals are manufactured from solid rings, split rings and segmented rings for use in both liquid and dry-running applications in the aerospace, nuclear, petrochemical and general marine industries.

Figure 2. Carbon-graphite seals are self lubricating, resistant to chemical corrosion, and capable of running at temperatures up to 538°C.
Carbon-Graphite Seals

Seal materials require high strength and a relatively high modulus of elasticity to withstand deformation at the interface. Carbon-graphite seal materials provide the strength and rigidity which are especially important in high pressure, zero leakage mechanical end-face seals. High thermal conductivity is essential in removing heat from the interface.

Seal wear is a result of adhesive wear, chemical wear, erosive wear and sometimes radioactive wear. Carbon-graphite is inert to most chemical reagents so it survives where other materials fail. However, chemical wear is evidenced in certain strong oxidising environments or where the additives are attacked by specific oxidising reagents.
Impregnation of Carbon-Graphite Seals

Impregnation of carbon seals can be done with a variety of materials to control permeability. In addition to thermoset resins, other types of impregnants include thermoplastics, metals, and inorganic salts or glasses. The temperature limit of the impregnant places an upper limit on the operating temperature of the carbon parts.

Metals such as antimony, silver, copper, nickel, and babbitt can improve the strength, thermal conductivity, and tribological characteristics of the materials. Impregnants made of inorganic salts usually phosphate or borate - and glasses are used in high temperature applications. Carbons impregnated with soluble salts must be handled carefully to avoid exudation, especially under humid conditions, but loss of impregnant rarely affects any physical property of a seal other than permeability.
Blistering

Blistering is a critical concern with carbon seal materials. Strangely, the reason why blistering occasionally occurs is not clear. One of the most popular explanations is that a certain amount of fluid becomes absorbed in the carbon substrate and expands due to frictional heat, creating a subsurface pressure and eventual crater in the seal face. Another theory is that softer mating materials can tend to tear pieces from the carbon-graphite in the presence of heavy hydrocarbons. Blistering is most often found in applications involving hydrocarbons or cyclical temperature service such as air conditioning compressors. In some cases, the use of silicon carbide as a mating surface will reduce or even eliminate a blistering problem, possibly because of its high thermal conductivity and hardness.
Carbon-Graphite Seals in Aerospace Applications

Carbon-graphite and graphites are excellent materials for aircraft turbine engine mainshaft seals. The mainshaft in a turbine engine rotates at very high speeds and operates in an environment of changing high temperature conditions. Mainshaft bearing compartment seals are used to protect rotor support bearings from hot gases flowing through the engine and to prevent the loss of lubricant in the bearing compartments.
Materials Selection and Suitability

Loads, speeds, temperatures, mating materials, cost constraints and projected volume are critical factors to be kept in mind when selecting materials. Scores of base carbon materials are available with hundreds of modifications that can be customised for specific designs and environments. General service carbon-graphites are usable up to 260°C while special grades are available that provide resistance up to 538°C. Special carbons and impregnants are used for seal applications in the 260-538°C temperature range.

These materials are chemically inert, temperature resistant, lightweight, resilient, dimensionally stable, and impermeable to gases and liquids. They can be moulded to size or machined to close tolerances, impregnated, plated, vulcanised to rubber, and cemented or shrunk into housings or retainers - a truly versatile set of materials.

Abstract

The objectives of this research are to synthesize and characterize nanomaterials of controlled size and shape so that they can be used in several different applications.  Nanophase metal particles have attracted a great deal of interest and have found applications in different fields such as catalysis, optical, microelectronic and magnetic devices and biological diagnostic probes due to their special properties that differ markedly from those of bulk materials.  Noble metal nanophases such as Ag, Pt, Pd etc., are the focus of catalytic applications because these nanoparticles can be hybridized with other substrates for enhanced catalytic functions.  A one-step microwave-assisted interface reaction using dodecylthiol and ethylene glycol was developed for the direct synthesis of hexagonally arranged spherical silver nanoparticles of about 10 nm.  Microwave-assisted solvothermal methods using ethylene glycol, ethanol or methanol as reducing agents were found to be useful for Pt and Pd nanoparticle synthesis at low temperatures.  Under conventional hydrothermal conditions, some nanowires such as Te were synthesized with the assistance of biomolecules.  Nanowires and nanorods of metals are expected to be useful as interconnects in electronic devices with super-functions.  Nanowires of Pt were grown in the nanochannels of mesoporous materials such as SBA-15, which served as a template.  Porous alumina membrane was used as a template for the growth of oriented SBA-15 nanorods with the nanoporous channels parallel to the channels of alumina and this hybridized material is expected to find super-function in nanowire fabrication and bio-molecule separations.  This review paper shows that conventional and microwave-assisted hydro- or solvothermal methods are eminently suited for the synthesis of nanomaterials of controlled size and shape under environmentally benign conditions.
Keywords

Nanowires, Nanophases, Solvothermal process, Hydrothermal Process, Biomolecules
Introduction

During the last two decades, there has been an enormous interest in nanostructures due to their conspicuous physical-chemical properties that differ markedly from those of bulk materials[1].  Various methods, such as hydrothermal and solvothermal routes[2], surfactant-assisted approach[3], have been utilized for the synthesis of nanomaterials.  Most physical and chemical properties of these nanomaterials are sensitively dependent on their size and shape, so materials scientists are still focusing on developing simple and effective methods for the fabrication of nanomaterials of controlled size and morphology[4].

Since metal nanoparticles have various applications, the synthesis of metal nanoparticles has attracted much attention especially in the last decade[5].  A variety of techniques have been developed to synthesize metal nanoparticles, including chemical reduction using a number of chemical reductants including NaBH4, N2H4, NH2OH, ethanol, ethylene glycol and N,N-dimethyformamide (DMF)[6-10], aerosol technique[11], electrochemical or sonochemical deposition[12, 13], photochemical reduction[14], and laser irradiation technique[15].  Because of the size-dependent properties, many physical, chemical and electrochemical methods have been employed to get the metal nanoparticles with uniform size, such as NaBH4-reduction approach resulting in the thiol-capped 1.8-3.5 nm diameter silver nanoparticles and alcohol reduction of fatty acid silver salts under microwave irradiation[16, 17].  The assembly of uniform nanoparticles into well-defined two- and three-dimensional (2D and 3D) superlattices is critically important to chemical, optical, magnetic and electronic nanodevices and would bring possibilities to brand-new properties and applications that result from the spatial orientation and arrangement of the nanocrystals[18].  Therefore, several approaches, such as self-assembly[19], Langmuir-Blodgett (LB) techniques[7], and electrophoretic deposition method[20] have been used in order to obtain self-organized lattices of metal, oxide and chalcogenide nanoparticles including silver[11], gold[21], cobalt[22], indium[23], α-Fe2O3[24], cobalt oxide[25], BaTiO3[26], CdS[27], CdSe[28], and Ag2S[29] nanoparticle arrays.

Besides the uniform and assembled nanoparticles, one-dimensional (1D) nanostructures, such as nanorods and nanowires, are also of particular interest not only because of their great potential for testing and understanding fundamental concepts but also because of their wide applications as interconnects in electronic devices with super-functions[30].  The synthesis of 1D nanostructures and guiding these nanometer-scaled building blocks to ordered superstructures would offer great opportunities to investigate the size- and dimensionality-dependent properties of these materials and could lead to the construction of nanoscale devices[31].  Until now, great progress has been made in the shape control of nanomaterials and a range of different 1D nanostructures have been fabricated by various techniques, such as Vapor-Liquid-Solid (VLS) growth mechanism[32], micro-emulsion method[3], hydrothermal (or solvothermal) technique[2] and template methods[33].  Among the various methods, hard template method is an effective method to obtain the nanostructures with low dimensionality.  Porous alumina membrane and mesoporous materials such as SBA-15 are two of the most used templates.  Nanowires of Ag, Pt, and Au were grown in the nanochannels of SBA-15[34], and many other nanorods arrays have been obtained by porous alumina membrane[35].  However, the pore size of the alumina membrane is from dozens of nanometers to several hundred nanometers and the SBA-15 is usually in powder form or as membranes with its channels parallel to plane of the substrate, which limits their applications in nanodevice fabrication[36].  Combining these two templates by introducing SBA-15 into alumina membrane channels is expected to find super-function in nanowire fabrication and bio-molecule separations.

This review paper shows that conventional and microwave-assisted hydro- or solvothermal methods are highly suited for the synthesis of nanomaterials of controlled size and shape under environmentally benign conditions for several different applications.
Experimental
Microwave-Assisted Solvothermal Synthesis of Metal Nanoparticles

For the hexagonally arranged spherical silver nanoparticles, 0.15 g AgNO3 was added in a Teflon vessel of a double-walled digestion vessel used in MARS-5 system.  Then 10 ml toluene, 1 ml dodecylthiol and 4 ml ethylene glycol were added into the vessel in order.  After sealing, the vessel was treated at 160°C for 3 hours using a microwave digestion system, MARS-5 (CEM Corp.).  After cooling to room temperature, the product was collected and an interface between two layers was found to be full of black product.

Pt and Pd nanoparticles were synthesized by microwave-assisted solvothermal method.  PVP with an average molecular weight of 40K was used as a capping agent in all the experiments.  Dihydrogen hexachloroplatinate (IV), and palladium (II) 2,4-pentanedionate were used as metal precursors.  PVP was dissolved in methanol or ethanol and then the metal salts were added.  The reactants were heated for 60 min at 90°C when methanol was used as a reducing agent and at 120°C when ethanol was used as a reducing agent for 60 min under microwave irradiation.
Biomolecule-Assisted Hydrothermal Synthesis For Te Nanowires[37]

For the elemental tellurium nanowires, 0.15g H2TeO4·2H2O was mixed with 0.075 g alginic acid in 10 ml distilled water in a Telfon-lined stainless autoclave.  After sealing, the autoclave was heated to 150°C and kept for 15 hours.  After cooling to room temperature, the solid product was collected by centrifugation at 2000 rpm for ~10 min and washed with distilled water and alcohol several times, followed by drying in air at room temperature.
Sol-Gel Method for The Growth of SBA-15 Nanorods Array Inside Porous Alumina Membrane[38]

In the synthesis of SBA-15 nanowire arrays inside porous alumina membrane, a sol solution was made by dissolving 1 g Pluronic P123 (PEO20PO70EO20, Mav=5800, Aldrich) in 5 g ethanol and 0.2 g 2 M HCl solution and mixing with 2.08 g tetraethyl orthosilicate (TEOS, 98%, Aldrich).  Then a simple piece of porous alumina membrane was put into the sol solution.  After the sol solution was left at room temperature (about 25°C) for 20 h to make sol change to gel, some amount of liquid paraffin wax with thickness of 1 mm was poured onto the gel and then it was kept at 60°C for 20 h.  Then, liquid paraffin was removed and the sample was calcined in the alumina membrane at 540°C for 6 h.
Nanowires of Pt Inside SBA-15

First, SBA-15 was prepared and treated with H2PtCl6 followed by H2 reduction at 400°C in order to prepare nanowires of Pt inside SBA-15.  The SBA-15 was then dissolved in dilute HF solution to recover Pt nanowries.
Characterization

The morphology, crystallinity, and size of products have been determined by transmission electron microscopy (TEM) and scanning electron microscopy (SEM).  Selected area electron diffraction (SAED) was used to identify the crystalline phases.  TEM was carried out with a Philips 420 transmission electron microscope operated at 120 kV and SEM was carried out with a Hitachi S-3500N scanning electron microscope.
Results and Discussion
Microwave-Assisted Solvothermal Synthesis of Metal Nanoparticles

Due to the spatial orientation and arrangement of the nanocrystals, 2D and/or 3D nanoparticles superlattices would bring possibilities to brand-new properties   and   applications, which make their syntheses be a focusing area in the current research field [18-20].  For the assembly of uniform Ag nanocrystals, the presynthesis of uniform nanoparticles or precursors is usually required followed by the organization process by surfactants or ligands.  The development of a simple and direct method for the fabrication of such crystals is a major challenge for future research.  Herein we report a general and one-step microwave-assisted interface-reaction for the synthesis and assembly of monodispersed silver nanoparticles.  By using dodecylthiol as directing reagent and ethylene glycol as reducing agent, hexagonally arranged spherical silver nanoparticles can be obtained by a one-step interface-reaction under microwave-assisted solvothermal conditions without the requirement of the pre-synthesis of uniform silver nanoparticles or special precursors and the technique of size-selective precipitation.  In the synthetic system, ethylene glycol and toluene form two layers with an interface where the thiol group of dodecylthiol might react with silver ions to form an inorganic-organic complex, which is reduced to elemental silver by ethylene glycol under microwave-solvothermal conditions. After the reaction, a black thin layer of silver nanoparticles is found at the interface and the formed silver nanoparticles automatically compact-pack to form ordered superstructures.  Figure 1 shows the TEM images of the as-prepared sample, from which it can be clearly seen that the sample consists of a hexagonal-like ordered superstructure of monodispersed silver nanoparticles.  Figure 1a displays a TEM image with low magnification, clearly showing that the two-dimensional (2D) hexagonal superlattice is the typical structure of the as-prepared silver sample.  A TEM image of the silver sample with high magnification (Figure 1b) displays clearly that these nanoparticles are monodispersed with an average diameter of ~ 10 nm and the inter-particle spacings are calculated to be about 2 nm.  Figure 1c shows its Fourier transform power spectrum.  It displays ordered hexagonal-like spot arrays, which confirms the formation of the hexagonally arranged silver superlattice.  The SAED pattern of the sample, showed in Figure 1d, exhibits polycrystalline diffraction rings, which can be indexed as cubic-phase metal silver, indicating that these nanoparticles are crystalline metal silver.

AZojomo - The "AZo Journal of Materials Online" TEM images, Fourier transform power spectrum and SAED pattern of the synthesized silver sample under microwave-assisted solvothermal conditions

AZojomo - The "AZo Journal of Materials Online" TEM images, Fourier transform power spectrum and SAED pattern of the synthesized silver sample under microwave-assisted solvothermal conditions

AZojomo - The "AZo Journal of Materials Online" TEM images, Fourier transform power spectrum and SAED pattern of the synthesized silver sample under microwave-assisted solvothermal conditions

AZojomo - The "AZo Journal of Materials Online" TEM images, Fourier transform power spectrum and SAED pattern of the synthesized silver sample under microwave-assisted solvothermal conditions

Figure 1. TEM images, Fourier transform power spectrum and SAED pattern of the synthesized silver sample under microwave-assisted solvothermal conditions.

Figure 2 shows the TEM images of synthesized Pt and Pd nanoparticles with methanol or ethanol as reducing agents.  Using methanol as reducing agent, Pt nanoparticles were synthesized.  Figure 2a shows the morphology of Pt nanoparticles formed with the PVP to Pt(IV) ratio of 18 and the concentration of Pt(IV) at 0.9 mM at 90°C.  The particle size is about 3 nm.  Pd nanoparticles can be also synthesized by using methanol as reducing agent.  Figure 2b shows the TEM image of Pd nanoparticles formed at 90°C with the PVP to Pd(II) ratio of 1.8 and the concentration of Pd(II) at 9 mM and the particle size is about 10 nm.  Using ethanol as reducing agent, Pt nanoparticles of approximately 3 nm were also obtained.  Figure 2c shows the TEM image of Pt nanoparticles formed at 120°C with the PVP to Pt(IV) ratio of 18 and the concentration of Pt(IV) at 9 mM.  Pd nanoparticles of around 10 nm were also synthesized by using ethanol as reducing agent.  Figure 2d shows TEM image of Pd nanoparticles formed at 120°C with the PVP to Pd(II) ratio of 18 and the concentration of Pd(II) at 9 mM.  Thus, Pt and Pd nanoparticles were successfully synthesized using either methanol or  ethanol  as  reducing  agents with microwave-assisted solvothermal technique.  Synthesized Pt nanopartices are about 3 nm and Pd nanoparticles are around 10 nm.

AZojomo - The "AZo Journal of Materials Online" TEM images of the obtained Pt and Pd nanoparticles under microwave-assisted solvothermal conditions

AZojomo - The "AZo Journal of Materials Online" TEM images of the obtained Pt and Pd nanoparticles under microwave-assisted solvothermal conditions

AZojomo - The "AZo Journal of Materials Online" TEM images of the obtained Pt and Pd nanoparticles under microwave-assisted solvothermal conditions

AZojomo - The "AZo Journal of Materials Online" TEM images of the obtained Pt and Pd nanoparticles under microwave-assisted solvothermal conditions

Figure 2. TEM images of the obtained Pt and Pd nanoparticles under microwave-assisted solvothermal conditions.
Biomolecule-Assisted Hydrothermal Synthesis for Te Nanowires

Bio-molecules, as life’s basic building blocks, have special structures with typical sizes in the range of about 5 to 200 nm, which is almost the same length scale as those of nanomaterials[39].  These biomolecules would be of great importance in developing novel materials and recently they have been introduced into the synthesis of nanomaterials[40, 41].  Alginic acid, a straight-chain polyuronic acid extracted from macrocystis pyrifera (kelp)[42], has been extensively used in pharmacy and cosmetic materials and recently as biosorption agent of heavy metals, which might be expected to be useful in controlled synthesis of nanomaterials.  Elemental tellurium has a wide range of applications in various thermoelectronics, photoconductors and piezoelectronic devices and the availability of 1D Te nanostructure could bring forth new applications or enhance the performance of existing devices[43, 44].  Herein we report a mild bio-molecule-assisted hydrothermal method by using alginic acid as both reducing agent and directional template to obtain 1D Te nanowires from commercial H2TeO4 powders under conventional hydrothermal conditions.

The as-synthesized nanowires’ structure and growth direction were characterized by TEM along with selected area electron diffraction (SAED) pattern.  Figure 3 is the TEM image of the obtained tellurium sample, which clearly shows that the obtained crystallites have a wire-like morphology.  The diameters of the Te nanowires are not very uniform and  the  average  diameter  is  calculated  to be about 80 nm and lengths are up to tens of micrometers.  Figure 3b and its inset display a single Te nanowire and its SAED pattern, which indicates that the nanowire might have [001] directional preferred growth.  All of the results clearly show that by using alginic acid as the reducing agent elemental tellurium could be obtained under mild hydrothermal conditions and the formed tellurium nanocrystallites have one-dimensional wire-like morphology.

AZojomo - The "AZo Journal of Materials Online" TEM images and SAED pattern of the obtained Te nanowires under biomolecule-assisted hydrothermal conditions

Figure 3. TEM images and SAED pattern of the obtained Te nanowires under biomolecule-assisted hydrothermal conditions.
Sol-Gel Method For The Growth of SBA-15 Nanorods Array Inside Porous Alumina Membrane

Mesoporous materials are special nanomaterials with ordered uniform nanochannels and would have important applications in various fields such as separation, catalysis, adsorption, advanced nanomaterials, etc.[45, 46].  SBA-15 has a highly ordered 2D hexagonal structure with adjustable pore size from 3 to 30 nm and high hydrothermal and thermal stability[46] and is expected to be useful in the synthesis of ultrafine nanorod arrays.  However, so far SBA-15 is still in its powder form or as menbrane with channels normally lying in the plane of the substrate, which limits its applications[36].  As an effective template, porous alumina membranes have stimulated great interest for the growth of ordered 1D nanostructures within their pores and up to now many nanorod arrays have been synthesized using porous alumina membrane as growth-limiting template[35].  Compared with SBA-15, alumina membranes have vertical one-dimensional (1D) channel structures, but with the pore sizes in the range of dozens of nanometers to several hundreds of nanometers, which limits its applications in the fabrication of nanodevices.  To combine the advantages of the alumina membranes and SBA-15 to form a membrane with fine, vertical mesochannels of about a few nanometers in size is of much importance and would provide wider applications in nanodevice fabrication and more extensive applications in other fields such as separation of biomolecules.  Herein we report a simple method for the synthesis of SBA-15 nanorod arrays inside the alumina membrane.

AZojomo - The "AZo Journal of Materials Online" SEM images of the obtained alumina membrane with SBA-15 nanorods inside

AZojomo - The "AZo Journal of Materials Online" SEM images of the obtained alumina membrane with SBA-15 nanorods inside

Figure 4. SEM images of the obtained alumina membrane with SBA-15 nanorods inside.

Figure 4 shows the SEM images of the obtained product. The top view SEM image (shown in Figure 4a) of the product evidently shows that nanorods are grown inside the pores of alumina membrane.  The diameters of the nanorods are in the range from 200 to 250 nm.  Figure 4b is the side view SEM image of the product, displaying that a number of nanorods grew inside the hexagonally arranged pore arrays of alumina.  These results clearly confirm that ordered SBA-15 nanorod arrays have formed in the channels of alumina membrane.  The mesoporous structure of SBA-15 was shown by TEM images.  Figure 5 shows TEM images of the obtained product, clearly displaying that the nanorods inside the porous alumina membrane have parallelly arranged channels with periodic spacing of ~ 9 nm, which is the (100) spacing of SBA-15.  The pore diameters of SBA-15 nanorods are calculated to be about 6 nm, which is the typical pore size for SBA-15 mesostructures.

From the above results, it is obvious that the SBA-15 nanorod arrays with vertical mesochannels are successfully obtained in the pores of alumina membrane that are used as a template.

AZojomo - The "AZo Journal of Materials Online" TEM images of the obtained alumina membrane with SBA-15 nanorods inside

AZojomo - The "AZo Journal of Materials Online" TEM images of the obtained alumina membrane with SBA-15 nanorods inside

Figure 5. TEM images of the obtained alumina membrane with SBA-15 nanorods inside.
Mesoporous SBA-15 As A Hard Template for The Growth Of Pt Nanowires

Figure 6 shows Pt nanowires of about 6 nm in diameter and about 100 to 200 nm in length grown in SBA-15.  It is very difficult to get smooth and long nanowires using mesoporous materials because of the difficulty in filling the mesopores completely with metal ions.  However, better nanowires can be obtained using biomolecule-assisted soft template method as was demonstrated with Te nanowires (See Figure 3).

 

AZojomo - The "AZo Journal of Materials Online" TEM image of Pt nanowires using mesoporous SBA-15 as hard template

Figure 6. TEM image of Pt nanowires using mesoporous SBA-15 as hard template.
Conclusions

In this review paper, some nanomaterials with controlled size and shape are successfully synthesized under microwave-assisted solvothermal or biomolecule-assisted hydrothermal method.

Hexagonally-ordered uniform spherical Ag nanoparticles of about 10 nm were obtained by a one-step interface reaction using dodecylthiol and ethylene glycol under microwave-assisted solvothermal conditions.  Pt and Pd nanoparticles were also synthesized under microwave-assisted solvothermal conditions at low temperatures using ethanol or methanol as reducing agents.  Under biomolecule-assisted conventional hydrothermal conditions, nanowires of elemental Te were synthesized with alginic acid as reducing agent and morphology-directing agent.  SBA-15 nanorod arrays with oriented mesochannels were obtained by using porous alumina membrane as template and this new and efficient mold is expected to find super-function in nanowire fabrication.  Platinum nanowires can be grown using SBA-15 as a hard template but the wires appear to be of poor quality.  These results show that by the conventional and microwave-assisted hydro or solvothermal methods, nanomaterials with adjustable size and shape can be successfully synthesized.
Acknowledgements

This work was supported by the NSF MRSEC under grant number, DMR-0213623 and the Huck Institutes of the Life Sciences.  TEM work was performed in the electron microscopy facility of the Materials Research Institute at Penn State University.

Bone Cement Composites

During the last 5 years bone cement materials have grown in popularity and are very promising osteoconductive substitutes for bone graft. They are prepared like acrylic cements and contain a range of powders such as monocalcium phosphate, tricalcium phosphate and calcium carbonate, which is mixed in a solution of sodium phosphate. These cements are produced without polymerisation and the reaction is nearly non-exothermic. The final compounds are reported to have a strength of 10-100 MPa in compression while 1-10 MPa in tension, although very weak under shear forces. These composites are currently used in orthopaedics in the management of fractures. It has been suggested that these materials could improve the compressive strength of the vertebral bodies in osteoporosis. Injection of calcium phosphate cement has been shown to be feasible and it does improve their compressive strength.
Hydroxyapatite Composite Materials

Preparation of hydroxyapatite/ceramic composites through the addition of various ceramic reinforcements has been attempted, metal fibres, Si3N4 or hydroxyapatite whiskers, Al2O3 platelets and ZrO2 particles. In many cases, the composites could not be successfully prepared and, because of problems related to a poor densification the mechanical properties could not be improved.

Hydroxyapatite/metal and hydroxyapatite/polymer composites are two typical classes of materials, which have been examined for improving the toughness characteristics of synthetic hydroxyapatite. In both cases, a toughness improvement can be found, due to a crack-face bridging mechanism operated upon plastic stretching of metallic or polymeric ligaments. Zhang et al. proposed a toughened composite consisting of calcium hydroxyapatite dispersed with silver particles. This material was obtained by a conventional sintering method. It was reported that the toughness of these composites increased up to 2.45 MPa m1/2 upon loading the mixture, with (30 vol%) silver. The use of silver is not only for taking advantage of the ductility of silver in terms of fracture toughness, but also because silver is inert and has anti-bacterial properties. Attempts to supersede metal alloys by carbon-fibre reinforced plastics and by various composites to stabilize fractures have met with limited success. Although a new titanium metal core composite hip implant has been clinically assessed in Europe with promising results.
Biomimetic Hybrid Composites

The conventional way to synthesize an inorganic material-based composite is to subject a mixture of the constituent phases to heat treatment. This process is also common in the biomaterials production arena; however, it is conceptually far from the biomineralization process, which occurs in nature. The natural process produces fine hybrid structures, which are hardly reproducible by classic consolidation processes. Traditional sintering route is not directly applicable to produce ceramic/polymer composites because no polymer will stand at the densification temperature of any ceramic material. Hydroxyapatite/polyethylene composites have been obtained by loading the polymeric matrix with the inorganic filler. In recent years, several research groups have demonstrated the feasibility of in vitro techniques for the synthesis of biomimetic material structures.

The sophistication of the biomimetic route has not been paired yet and these techniques, so far, have not proved to be fully applicable for clinical applications although various companies started to produce a range of clinical products. It can be easily predicted that more and more dense hybrid materials will be introduced, opening a completely new perspective in biomaterials production and application methods.

A new alternative route, -based on an in situ polymerization process, carried out into an inorganic scaffold (with submicrometer-sized open porosity)-, has also been recently proposed. This method is an intermediate one between conventional sintering and biomineralization in vitro, because it still employs sintering for the preparation of the inorganic scaffold, but the subsequent hybridisation of the scaffold with organic phases is carried out through a chemical route. This method enables the synthesis of biomimetic (hybrid) inorganic/organic composites, while aiming at relatively complex structural designs; it is rather simple and easily reproducible process. A schematic of this efficient synthesis route is given in Figure 1.

Figure 1. Schematic representation of the in situ polymerisation synthesis route of new generation hybrid materials.

A common characteristic of natural biomaterials such as bone, nacre, sea urchin tooth and other tough hybrid materials in nature is the strong microscopic interaction between the inorganic and the organic phases. This characteristic allows the organic phase to act as a plastic energy-dissipating network, forming stretching (bridging) ligaments across the faces of a propagating crack in a nanoscale level. Such complexity has led to the common perception that, to mimic natural designs, in situ synthesis techniques should be adopted. Precipitation of calcium carbonate or hydroxyapatite into a polymeric matrix, for example, has been proposed as a novel synthetic route to biomimetic composites. Despite significant advances in understanding biological mineralization and developing new fabrication processes, the composites to date obtained by these methods are by far in embryonic stage for actual applications, due to their low structural performance.

The results of fracture tests carried out on two natural biomaterials, bovine femur and Japanese nacre (Crassostrea Nippona), in comparison with a synthetic hydroxyapatite/nylon-6 composite obtained by in situ polymerisation of ε-caprolactam infiltrated into a porous apatite scaffold showed that the high work of fracture achieved is about two orders of magnitude higher than that of monolithic hydroxyapatite, and it is due to stretching of protein or polymeric ligaments across the crack faces during fracture propagation.
Summary

Although the nanoscale modelling of synthetically manufactured hybrids and composites is still in infancy, mimicking natural microstructures while using strong synthetic molecules may lead to a new generation biomaterials, whose toughness characteristics will be comparable with the materials available in the nature. A formidable challenge remains on the optimisation of their morphology and bioactivity in these novel hybrid composites.

A complete set of references can be found by referring to the original paper.

Abstract

On the rear panel in Plasma Display Panel, barrier ribs were formed in stripe and matrix arrangements to maintain the discharge space between the two glass plates as well as to prevent electrical and optical cross talking between the adjacent cells.  Various barrier rib forming routes were proposed including screen printing, sandblasting, lift-off, rolling of green tape, and photosensitive paste.  In this study, a rear panel forming process by photosensitive paste was presented.  Photosensitive barrier rib pastes were prepared by incorporating glass powders, binder polymer, functional monomer, photoinitiator, additives, and solvent.  After optimization of the paste formula and photolithographic process, a barrier rib could be obtained with a high resolution of 150μm height and 30μm width, with a pitch of 150μm.
Keywords

Plasma Display Panel, Barrier Rib, Photolithography, Photosensitive Paste, Fine Pattern
Introduction

On the rear panel in Plasma Display Panel (PDP), barrier ribs were formed in stripe and matrix arrangements to maintain the discharge space between the two glass plates as well as to prevent electrical and optical cross talking between the adjacent cells, and also to provide the additional areas for phosphor coatings on their side walls, which contribute to the improvements in contrast as well as luminance of the device.  The width of the barrier rib used in a PDP is typically in the range 30-80μm and its height is 110 –140μm.  The distance between the barrier rib centers (subpixel pitch, pitch) is 150 and 420μm for 42 inch (107 cm) full-spec.  High Definition Television (HDTV), 1920×1080 pixels) and Video Graphic Adaptor (VGA), 852×480 pixels) grade PDPs, respectively.  Ribs with finer dimensions have been required for devices of higher resolution.

Various barrier rib forming routes were proposed and examined including screen printing [1], sandblasting [2], lift-off, rolling of green tape [3], and photosensitive paste [4-6] as shown in Figure 1.  Among the routes, sandblasting and photosensitive paste were currently used in most PDP- manufacturing lines.

AZoJomo - The AZO Journal of Materials Online - Schematic figure of various barrier rib forming processes

Figure 1. Schematic figure of various barrier rib forming processes.

In photosensitive paste process, a glass substrate is coated with a photosensitive glass paste and the patterning is carried out by a conventional photolithographic process.  Conventional photosensitive paste process was, however, not yet in practical use as a rib forming though photosensitive silver paste has been widely used for the formation of sustain, bus, and address electrodes [7, 8].

By the conventional photosensitive paste process, a thickness of the pattern attained by one time exposure is as thin as 30μm as shown in Figure 2.  Therefore, it has a big shortcoming to repeat the application of paste, drying and exposure for 4-10 times to form a rib of a required height, i.e.200μm (about 140μm after firing).

AZOJomo - THe AZO Journal of Materials Online - Problems of barrier rib forming by conventional photosensitive paste process

Figure 2. Problems of barrier rib forming by conventional photosensitive paste process.

In the present work, this shortcoming of repeating the exposure and development was solved by the examinations of fabrication process, namely paste compositions and glass powders design.  The chemical composition, particle size, and distribution of glass powders and firing conditions were investigated.  As a result, photosensitive paste process was established which could form the required rib pattern by one time application, drying, exposure, and development as shown in Figure 3 [9, 10].

AZoJomo - The AZO Journal of Materials Online - New barrier rib forming process by present photosensitive paste

Figure 3. New barrier rib forming process by present photosensitive paste.

Photosensitive rib pastes were prepared by incorporating the glass powders, binder polymer, functional monomer, photo initiator, additives, and solvent.  The effects of the paste components and photolithographic processes on rib forming were examined.  The effects of the processing parameters such as the exposure energy and the gap between the coated film and photo mask, and the development conditions on the height, width, pitch, and cross section of the ribs were investigated.
Experimental
Photosensitive Paste Preparation

The photosensitive paste process is schematically shown in Figure 4.  Photosensitive pastes are composed of inorganic glass powders and photosensitive organic components including acrylic binder polymer, ultraviolet (UV) curable functional monomer, photo initiator, plasticizer, additives, and solvent.

AZoJomo - The AZO Journal of Materials Online - Schematic image of photosensitive paste process

Figure 4. Schematic image of photosensitive paste process.

Inorganic glass powder was produced by glass makers and its chemical composition was within the B2O3-SiO2-Al2O3-Li2O-BaO-MgO system.  The average particle sizes of the glass powders, measured using laser diffraction particle size analysis (Microtrack) were 1.0-4.0μm.

Pastes for rib forming were prepared by kneading the glass powders and organic constituents, i.e., solvent, polymer binder, monomer, plasticizer, and other additives.

First, binder polymer was dissolved in γ-BL (butyrolactone) solvent with mechanical stirrer to the extent which did not have any gel or coaglulums.  Monomers, photo initiator, plasticizer, and additives were added to this solution, and the resulting mixture was stirred.  To this mixture, the slurry and glass powder were added and mixed with a mechanical stirrer.  Total slurry of desired formulation was then kneaded by the three roll mill (Exact Co., Germany) for 5 -10 times to achieve the uniform dispersion and desired viscosity.
Printing

Photosensitive paste was screen printed over the whole surface of the glass substrate of high strain point (Asahi Glass PD-200, 125 mm square and 2.8 mm thickness).  The stainless steel screen of 325 mesh (25μm in opening) was used.  Printing and drying were repeated over 10 times so that the thickness reached 200μm.  After leveling at room temperature for 5-10 min., the coated film on glass substrate was dried at 80˚C to remove the solvent.
Barrier Rib Patterning by Photolithographic Method

Using glass photo masks with three different pitches (360, 230, and 130μm) of chromium stripes, the coated substrate was exposed to UV irradiation (wave length: 365, 405, and 436 nm).  Light source was ultra high voltage mercury lamp and the amount of exposure is varied from 200-700 mJ/cm2.

Exposed substrates are immersed into a weak alkali aqueous solution of 0.1-0.5 % for 10-30 sec. to dissolve the unexposed portion, and then dried at 80˚C to obtain the barrier rib formed substrate.  Barrier rib formed substrate was fired in an air atmosphere using a box type furnace.  Temperature was heated up to 580˚C at a rate of 80 -200˚C/h and kept at that temperature for 15 min to sinter.

Barrier rib patterns without any binder polymer and other organic components were obtained through this firing process.

Cross section, top view, and the defects such as disconnection of fired patterns were examined using scanning electron microscopy (SEM) and optical microscope.
Results and Discussion

The patterning of barrier rib is dependent on many factors and parameters involved in both the formulation of photosensitive pastes and photolithographic processes.  One of the important factors controlling the photosensitivity and photo polymerization of barrier rib paste is the selection of UV curable monomer/polymer system and photoinitiator.  Perfect photo polymerization throughout the bottom part of the barrier rib is a key factor to obtain the good patterning without disconnection and curving.

Figure 5 shows the SEM images of cross section and whole view of barrier ribs.  Each of the ribs with 360, 230 and 130μm pitches was formed by one time exposure and development.  In order to form a rib pattern of 200μm height, the UV scattering by photosensitive paste should be avoided.  The ribs with high resolution (XGA) of 150μm height, 30μm width, and 150μm pitch after firing could be formed by the optimization of paste formula and photolithographic process as shown in Figure 6.  Any defect such as the disconnection, peeling, and curving of ribs was not detected.

AZoJomo - The AZO Journal of Materials Online - SEM image of rib pattern of 360, 230 and 130μm pitches by photosensitive process and their applications.

Figure 5. SEM image of rib pattern of 360, 230 and 130μm pitches by photosensitive process and their applications.

AZOJomo - The AZO Journal of Materials Online - SEM image of rib pattern of 150μm height, 30μm width, and 150μm pitch by photosensitive paste

Figure 6. SEM image of rib pattern of 150μm height, 30μm width, and 150μm pitch by photosensitive paste.

In addition, two different types of the cross section of ribs could be formed by selecting the paste formula and exposure energy to obtain the uniform phosphor layer thickness for high brightness as shown in Figure 7.

AZOJomo - The AZO Journal of Materials Online - SEM image of rib pattern of two different cross sections

Figure 7. SEM image of rib pattern of two different cross sections.
Conclusions

The barrier ribs with high resolution (XGA) of 150μm height, 30μm width, and 150μm pitch after firing were achieved by the optimization of paste formulation and photolithographic process.  The photosensitive paste process is especially suitable for high resolution plasma display panel.

Ferrous Metals

The major products of the ferrous industry are cast iron, low and medium alloyed steels and specialty steels such as tool steels and stainless steels. In the production processes used to manufacture these end products other important materials encountered are refractories, sinters and slags, and of course the starting raw materials iron ore together with limestone and coke.
Ferro Alloy Additions for High Alloy Steels

Finally ferro alloys are employed in the high alloy steel making processes as a way of adding elements such as chromium nickel, molybdenum and silicon to the steels thereby completing an extremely varied spectrum of materials to be analyzed in the ferrous industry.
Analysis Requirements for Low Alloy Steels

For the bulk produced low alloy steels, where trace and tramp elements exerts significant effects on the strength and ductility, low detection limits and high accuracy are needed coupled with a fast analysis time.
Analysis Requirements for Austenitic Stainless Steels

For austenitic stainless steels where costly and often strategic alloying elements are added in relatively large quantities, accuracy and high element concentration level is demanded.
Analysis Using MagiX FAST from PANalytical

With MagiX FAST from PANalytical, measuring results can be achieved for wide different metallic matrices indicating detection limits below 10 ppm for many elements in measuring times of less than one minute.
Non-Ferrous Metals

The non-ferrous industries deal with a similar cross-section of material types in the production of metals such as copper, nickel, lead and zinc, but in these industries much more time and effort is spent in analyzing exploratory core body samples since these metals are generally present in much lower concentrations in the earths crust than is iron and more widely dispersed.
Analysis Requirements for Aluminium and Copper Alloys

Aluminum and copper are widely used in high electrical conductibility applications where the presence of some elements at even a trace level can give rise to unacceptable high power losses in electrical transmission lines. For aluminum an added concern is the fact the anodizing characteristics of the metal are strongly influenced by the presence of just a few ppms of certain elements. Thus for both, copper and aluminum, the most relevant analytical data concern is the detection limits achievable with a spectrometer.
PANalytical Analysis Solutions

The unsurpassed sensitivity, reproducibility and stability of Axios-Advanced make it possible to achieve these low limits of detection. The simultaneous PANalytical MagiX FAST X-ray spectrometer can analyze up to 28 elements simultaneously in less than one minute with a high degree of precision.
The PANalytical EDXRF Bench-Top Systems

The PANalytical EDXRF bench-top systems MiniMate and MiniPal are ideally suited for single element analysis and as back-up systems of larger laboratory equipment.
Potflux X-Ray Diffraction Systems

X-ray diffraction is the aluminium industry’s standard method for potflux process control. The PANalytical CubiX PRO Potflux X-ray diffraction system allows simultaneous measurement of the phases present in the aluminium bath, as well as the calcium (Ca) content through an integrated Ca XRF channel in the diffractometer. Ultrafast pneumatic linear sample introduction means high throughput for the customer - as many as 60 samples per hour can be analyzed with the CubiX PRO Potflux.
Precious Metals

For the analyses of precious metals such as Au, Ag, Pt, Rh, Ru, Ir, Pd, Cu, Zn, Ni, Co, etc often XRF systems are used since a vital part of working with precious metals has to be the ability to determine accurately the composition of material at all times. It is important to be able to monitor the composition of work-in-progress both to control and minimize metal losses and able to effect regular inventory checks.
Advantages of X-Ray Fluorescence Over Traditional Techniques

Traditional methods like the Touchstone test, cupellation, fire-assay, nitric acid tests and also other analytical methods are destructive and often time consuming and hazardous. Also the preparation of samples takes longer.
Advantages of X-Ray Fluorescence Analysis

XRF as a non-destructive analysis technique can identify and determine the presence of various elements such as gold, platinum, silver, and other precious metals present in solid, powdered and liquid samples often within a matter of minutes. Sample preparation is quick and easy and solid samples as small as 3 centimetres can also be analyzed. Essential advantages of using XRF are the high reproducibility or precision and the high accuracy of this analytical method.
PANalytical XRF Solutions For Precious Metals

The Epsilon 5 is a high-performance EDXRF spectrometer specifically designed to achieve the lowest possible detection limits for heavy elements such as Cd, Pb, As, precious metals and the lanthanides.

AZoM - Metals, Ceramics, Polymer and Composites : Analysis of Ferrous Alloys, Non-Ferrous Alloys, Precious Metals and Metal Coatings Using X-Ray Systems – The Epsilon 5 from PANalytical

Figure 1. The Epsilon 5 EDXRF Spectrometer from PANalytical.

PANalytical offers a comprehensive range of products starting with the portable bench-top systems MiniMate and MiniPal through the medium-sized instruments up to the fastest and most sensitive spectrometer available today, the Axios-Advanced.
Metal Coating

Steel producers need an excellent quality control of coatings during their entire processes in order to avoid production loss, waste of coating material and achieve fastest possible control over their process. Car manufacturers need to quickly analyze different coating weights in order to keep quality at highest standards.
Advantages of X-Ray Fluorescence for Analysis of Metal Coatings

The analysis of metal coating is an ideal application for XRF because X-rays achieve higher penetration depths than comparable techniques. Coatings represent a variety of metallic mixtures and compound typically used to strengthen certain features of a product to make it better fit its purpose. And the recipes for these mixtures solely depend on what the user will do with the coated product.
Typical Metal Coatings

Typical metals used for coatings among others are titanium, aluminum, zinc, zirconium, chromates and phosphate for example for steel pre-treatment. TiN e.g. works well for machining iron-based materials. CrC with its high temperature oxidation resistance, is used in die casting while Tungsten carbide/carbon (WC/C) is designed to coat and protect highly-loaded precision components, gears and gear drives, engine components and hydraulic pumps and compressors.
Common Coatings Scenarios Where XRF Analysis is Used

Common coating applications where XRF is used as standard analysis tool include for example conductive coatings on plastic. Other typical applications are metalized food packaging and solar power cells using thing coatings of a metal alloy on a polymer subtrate. XRF systems are used for the precise non-destructive measurement of coating thickness as well as for alloy coatings and in the incoming material inspection.
PANalytical XRF Tools For Coating Analysis

The PANalytical EDXRF bench-top systems are typical XRF tools used in metal coating analysis. For more complex task we offer a comprehensive range of WDXRF spectrometers. The FP-Multi option of the SuperQ software package uses fundamental parameters to determine the thickness and composition of coatings and surface or sub-surface layers on samples such as metals and semiconductor wafers.

The carbides constitute a particularly interesting family of compounds in that they were the first man-made refractories and, unlike oxides and silicates, are extremely uncommon in nature. Iron carbide, Fe₃C, and later TiC and WC were identified in and extracted from steels in the mid 1800’s. By 1900, the French chemist Moisson had synthesized a number of refractory carbides in his new arc furnace and had studied their properties.

Characteristically, most of these carbides have high hardness, good electrical and thermal conductivity, and high stability. These properties account for the principal applications: structures resistant to chemical reaction, uses in which wear resistance is of major importance, and high- temperature radiant-energy sources. The brittleness of carbides, however, has prevented their use as single-phase materials in highly stressed structural applications and has led to the development of metal-bonded composites.

The carbides vary widely in their chemical inertness, particularly with respect to their attack by oxidizing atmospheres. At sufficiently high temperatures all are attacked quite rapidly, although some are more resistant than metals of comparable melting point.

Until recently, adequate attention had not been given to the importance of structure in determining properties. Not only microstructure but also lattice-defect structure and dislocation substructure are important. Especially among interstitial compounds, details of structure and chemical composition exert an important influence on mechanical and physical properties. Consequently, few properties of the carbides are known with precision and their future development is, as yet, unbounded.

Carbides may be grouped according to the periodic classification of the metal constituents. The carbides of boron, silicon, and the transition metals (including the rare-earth and actinide series) are of greatest interest and utility.

The manufacture of carbides involves two groups of processes: those resulting in synthesis of the compound and those involving the purification and consolidation of the compound into an integral body. In some preparative techniques, such as vapor plating, these groups may be combined.

The synthesis techniques may be subdivided as shown in Table 1. A specific example of each is shown. The exact nature of the reaction, however, seldom is known in detail and is strongly dependent on temperature and other variables. Although the reactions usually are classified as noted above, it must be remarked that even in solid-solid or solid-liquid reactions the presence of a vapor phase is often critical for completion or rapid rate of reaction.

The processes of greatest present commercial importance are the solid-state carburization of the metal and the carburization of the oxide, but some of the other reactions are inherently capable of yielding products of higher purity and, hence, are of greater interest for research purposes and special applications.

The principal consolidation techniques have been sintering, hot pressing, and surface deposition. Most sintering of transition-metal carbides has been carried out in the presence of a liquid metallic phase. Relatively little has been done on the sintering of carbides without the use of additives which form a liquid phase.

The usual procedure for sintering unadulterated transition-metal carbides is to fire the pressed compact in the best available atmosphere, well gathered with raw carbide or reactive metal, at the highest feasible temperature. Final sintering at extremely high temperatures is accomplished by passing an electric current directly through the presintered bar.

The very high sintering temperatures required result in some purification by volatilization of impurities. Temporary liquid-phase formation with small metal additions has sometimes been used, the metallic phase disappearing at the end of the sintering operation by either solid-solution formation or volatilization.

Little has been done on solid-state sintering of carbides with "activating additions," as has been extensively employed in oxide and metal sintering. Hot pressing allows densification of unadulterated carbides at lower temperatures than does solid-state sintering. A variety of surface-deposition methods has been examined in recent years for providing wear-resistant, corrosion-resistant, or special purpose coatings of carbides on different substrates. Among these are the vapor-deposition techniques, flame spraying, and plasma jets.

It now appears, however, that electronic as well as size factors must be considered and the carbides more properly may be regarded as metallic lattices stabilized by electron transfer from the carbon atom. The viewpoint that an optimum number of bonding electrons exists for a given coordination number finds support in the extraordinary stability of these carbides and in several other observations: true stoichiometry is some- times energetically unfavorable [e.g., TaC and TiC-WC solid solutions, melting-point maxima are observed in some quasi-binary systems half of the tetrahedral sites formed by cubic close packing of the other type. Various analogous hexagonal and rhombohedral arrays of tetra-hedrally coordinated carbon and silicon are all known as a-SiC. Although SiC itself is practically stoichiometric, some A1₄C₃, which in some respects is similar to SiC, is soluble in it.

Boron carbide, B₄C, is similar to the rhombohedral modification of boron in having icosahedra of 12 boron atoms centered on the lattice points of a deformed face-centered cubic lattice and bonded together along the diagonals of the lattice. Centered on each octahedral interstice, making a rhombohedral NaCl lattice, are three carbon atoms in a chain. The end carbon atoms in every chain are bonded to three boron atoms, one in each of the three closest icosahedra.

The resulting structure has exceptional hardness and stability, although the central carbon atom can be removed easily by neutron irradiation and can be replaced by boron to yield B₁₃C₂ at high temperatures. Both titanium and silicon dissolve in the structure in small amounts. The lattice is capable of retaining more helium, resulting from neutron irradiation, than can pure boron.

Other carbides of interest include the strongly ionic MC₂ structures (Group II, rare-earth, and actinide series) and Be₂C.

The most stable transition-metal carbides are composed of close-packed layers of metal atoms stacked together in several arrangements. Carbon sites are either trigonal prismatic, octahedral, or tetrahedral holes formed between the metal layers: WC; W₂C ; MoC; UC₂; ,S-SiC; a-SiC; and Be₂C.

In the more complex transition-metal carbides, carbon usually occupies sites similar to those in the structures described above. In Cr₃C₂, carbon occupies trigonal prismatic holes, as in WC, except that adjacent sites are occupied to form parallel zigzag chains of carbon through the structure. The nearest neighbors of carbon in the MeC "eta " carbides are identical to those in W₂C, but the structure as a whole is considerably more complex, with 112 atoms in the unit cell. In (Cr, Fe, W)₂₃Ce, the carbon atoms lie among 8 metal atoms instead of 6, but the M-C bond distances are the same.

Two of the most important carbides, SiC and B₄C, have strongly directed, covalent bonds and are not interstitial compounds, although structural similarities are evident.

Mechanical properties and stress-strain behavior were evaluated for several types of commercially fabricated aluminum matrix composites, containing up to 40% vol discontinuous SiC whisker, nodule, or particulate reinforcement. The elastic modulus of the composites was found to be isotropic, to be independent of type of reinforcement, and to be controlled solely by the volume percentage of SiC reinforcement present. The yield/tensile strengths and ductility were controlled primarily by the matrix alloy and temper condition. Type and orientation of reinforcement had some effect on the strengths of composites, but only for those in which the whisker reinforcement was highly oriented.

Ductility decreased with increasing reinforcement content; however, the fracture strains observed were higher than those reported in the literature for this type of composite. This increase in fracture strain was probably attributable to cleaner matrix powder, better mixing, and increased mechanical working during fabrication. Comparison of properties with conventional aluminum and titanium structural alloys showed that the properties of these low-cost, lightweight composites demonstrated very good potential for application to aerospace structures.

The majority of effort in aluminum matrix composites has been directed toward development of high performance composites, with very high strengths and module, for use in specialized aerospace applications.

However, there are a number of other applications in aircraft engines and aerospace structures where these very high properties may not be required, and where it could be cost effective to use other metal matrix composites. For example, cost-, weight-, and stiffness-critical components, such as engine static structures, do not require the very high directional properties available with composites reinforced with aligned continuous fibers. Replacement of such current aluminum, titanium, or steel structures by low cost composites offers the potential of significant weight and cost savings.

For these reasons, efforts were initiated to assess the potential of applying low cost aluminum matrix composites to these structures, using low-cost reinforcements and low-cost composite fabrication processes, including powder metallurgy, direct casting, and hot molding techniques.

Factors Influencing Modulus of Elasticity

The modulus of elasticity of 6061 Al matrix composites increased with increasing reinforcement content. This increase, however, is not linear, as in the case of composites with continuous fibers aligned in the testing direction. The modulus of the composites was below that expected from isostrain-type rule-of-mixtures behavior, and tended to approach an isostress-type hyperbolic function with reinforcement content, similar to that observed for transverse modulus behavior of continuous fiber composites.

The reinforcement content was the dominant factor in the improvement of modulus of elasticity in these SiC/Al composites. For a given reinforcement content, the modulus tended to be isotropic with nearly equal values obtained from tests in both the longitudinal and transverse directions. In addition, the modulus appeared to be independent of type of reinforcement, with modulus values being within 5% of the average value for all composites nested at any given reinforcement content, regardless of type of reinforcement.

The modulus of the composites was also independent of the matrix alloy. Heat treatment of the composites may have had a slight effect on modulus. The modulus of composite in the T6-temper appeared to be slightly lower than the modulus measured on composites in the as-fabricated F-temper. This reduction was slight (about 3 to 4%) and was not consistent among all the matrix alloys tested, and may have been due to scatter in the data.

Factors Influencing Strength

The factors influencing the yield and tensile strengths of SiC/Al composites are complex and interrelated, and the best way to evaluate this behavior is through isolation of variables and analysis of stress-strain curves and fracture behavior.

Effect of Al matrix alloy. The Al matrix used for the SiC/Al composites was the most important factor affecting yield strength and ultimate tensile strength of these SiC/Al composites. Tests showed that SiC/Al composites with higher strength aluminum matrix alloys, such as 2024/2124/7075 Al had higher strengths but lower ductilities.

Composites with a 6061 Al matrix showed good strength and higher ductility. Composites with a 5083 Al matrix failed in a brittle manner, with ultimate strength related to failure strain. The 5083 Al alloy is not heat-treatable and has been optimized to gain maximum properties by solid solution strengthening in the strain-hardened H-temper. The addition of the SiC reinforcement probably overstrained the lattice, and thus the alloy no longer had sufficient strain energy remaining to gain its potential strength and ductility.

While heat treatment had little, if any effect of the modulus of elasticity of the composites, it did affect the transition into plastic flow. Composites in the F-temper strained elastically and then passed into a normal decreasing-slope plastic flow.

Composites tested in the T6-temper exhibited a slightly greater amount of elastic strain, with the elastic proportional limit being increased from about 0.10 to 0.15% strain to about 0.15 to 0.25% but the greater influence was a steepening of the slope of the stress-strain curve at the inception of plastic flow, relative to that observed for composites in the F-temper. The inception of plastic flow was marked by a continuation of a slope that, while no longer elastic and starting to become plastic, approached that of the elastic portion. This slope decreased with increasing strain, until eventually reaching normal plastic flow leading to fracture at the ultimate tensile strength.

This increase in elastic proportional strain limit and steepening of the stress-strain curve were reflected by the higher yield and ultimate tensile strengths observed in the heat-treated composites. The increase in flow stress of composites with each heat-treatable matrix probably indicated the additive effects of dislocation interaction with both the natural alloy precipitates and the synthetic SiC reinforcement. The combination increased the lattice strain in the matrix, causing greater dislocation tangling and requiring higher flow stresses for deformation, resulting in the higher strengths observed.

Experiments showed that the yield and ultimate tensile strengths of the SiC/Al composites, with other parameters being constant, were primarily controlled by the intrinsic yield/tensile strengths of the matrix alloys. Also, the yield and ultimate tensile strengths of the composites, with 20% pct SiC reinforcement, were shown to be higher than those of the same heat treated matrix alloys without reinforcement. The largest increase in yield/tensile strengths over those of the unreinforced matrix alloy was achieved by the SiC/6061 Al composites.

Factors Influencing Ductility

Ductility of SiC/Al composites, as measured by strain to failure, is again a complex interaction of parameters. However, the prime factors affecting these properties are reinforcement content, matrix alloy, and orientation.

With increasing reinforcement content, the failure strain of the composites is reduced, and the stress-strain curves also reflect a change in the fracture mode. Preliminary tensile tests, conducted on wrought aluminum specimens with no SiC reinforcement, exhibited failure strains of about 15 pct, with a smooth 45 deg chisel-point shear fracture across the thickness of the specimen. There was also a contraction in the width of the specimen at the fracture plane.

Elevated Temperature Properties

Discontinuous SiC/Al composites continued to show an advantage over conventional aluminum alloys at elevated temperatures.

Specimens tested at temperatures of 149° to 204°C (300° to 400°F) exhibiting the same type of V-shaped, double shear lip transition fracture observed in tests at room temperature. Specimens tested at 260°C (500°F) showed a slight increase in plastic strain. While still transitional, the fracture showed more of a tendency for the formation of a more ductile, single shear lip and was basically the same as that observed at lower temperatures. Failure strain appeared to increase slightly at 315°C (600°F). The fracture showed a great deal of necking in both the width and thickness direction of the specimen, and all four surfaces of the fracture area necked in a ductile manner. This change in fracture behavior coincided with the marked drop in ultimate tensile strength observed at 315°C (600°F).

Application of SiC/Al Composites to Aircraft Engine and Aerospace Structures

Studies show that these low cost SiC/Al matrix composites demonstrated a good potential for application to aerospace structures and aircraft engine components. The composites are formable with normal aluminum metal-working techniques and equipment at warm working temperatures. They can also be made directly into structural shapes during fabrication.

These composites merit additional work to determine fatigue, long-term stability, and thermal cycle behavior to characterize more fully their properties and allow their consideration for structural design for a variety of aircraft and spacecraft applications.

The most significant aspect of these data was the increase in modulus over that of competitive aluminum alloys. At 20 vol pct reinforcement, the modulus of SiC/Al composites was about 50% above that of aluminum and approached that of titanium. This increase in modulus was achieved with a material having a density one-third less than that of titanium. Comparison of the properties of the various composites shows that the modulus/density ratio of 20 vol pct SiC/Al composites was about 50% greater than that of Al or Ti alloys, while at 30 vol pct SiC the advantage was increased to about 70% and at 40 vol pct SiC the modulus was almost double that of unreinforced Al or Ti structural alloys.

Studies were undertaken to evaluate the tensile behavior of low-cost discontinuous SiC/Al composites, containing SiC-whisker, -nodule, or -particulate reinforcement. The effects of reinforcement type, matrix alloy reinforcement content, and orientation were determined by analysis of stress-strain curves and by SEM examination. These investigations led to the following conclusions:

1. Discontinuous SiC/Al composites offer a 50 to 100% increase over the modulus of unreinforced aluminum and offer a modulus equivalent to that of titanium, but at a third less density. The SiC/Al composites had modulus/density ratios of up to almost twice those of titanium and aluminum structural alloys. The modulus of SiC/Al composites tended to be isotropic and was controlled by the amount of SiC reinforcement.
2. The yield and tensile strengths of SiC/Al composites demonstrated up to a 60% increase over those of the unreinforced matrix alloys. Yield and ultimate tensile strengths of the composites were controlled by the type and temper of the matrix alloy and by reinforcement content. In general, these properties were independent of the type of reinforcement.
3. Ductility of SiC/Al composites, as measured by strain to failure, was dependent upon reinforcement content and matrix alloy. Composites with ductile matrix alloys and lower reinforcement contents exhibited a ductile shear fracture with a 5 to 12% failure strain. As reinforcement content increased, the fracture progressed through a transition and became brittle, reaching a <1 to 2% failure strain, at higher reinforcement contents. The increase in ductility over that reported previously was probably attributable to cleaner matrix alloy powders, better mixing, and increased mechanical working.
4. A fine dimple network was observed in the fracture surfaces of composites with higher strains. At lower fracture strains, a coarser dimple network was observed. Composites failing in a brittle manner showed increasing amounts of cleavage fracture.
5. The SiC-whisker reinforcement was generally oriented in the extrusion direction. Composites with a higher degree of preferred orientation tended to have higher ultimate tensile strength in the direction of whisker orientation. Composites with a more random whisker orientation tended to be isotropic in strength.

The expected benefits of economical, high-performance civil-aircraft designs that are being considered for the future will be realized only through the development of light-weight, high-temperature composite materials for engine applications to reduce weight, fuel consumption, and direct operating costs.

A major effort underway in this area is the Advanced High Temperature Engine Materials Technology Program (HITEMP) of the National Aeronautics and Space Administration (NASA), which focuses on providing revolutionary high-temperature composite materials: to 425°C for polymer-matrix composites (PMCs); to 1250°C for metal-matrix / intermetallic-matrix composites (MMCs / IMCs); and to as high as 1650°C for ceramic-matrix composites (CMCs).

Composites promise benefits
Numerous conducted studies demonstrate that significant economic and performance benefits can be achieved if lightweight, high-temperature composite materials can reach technology readiness. Based on a preliminary design of a conceptual engine, however, material temperatures approaching 1650°C are anticipated for the turbine inlet, thus requiring extensive use of CMCs throughout the combustor, turbine, and exhaust nozzle.

One benefit of using CMCs is that they allow higher operating temperatures and thus greater combustion efficiency leading to reduced fuel consumption. Thanks to the low density of CMCs, compared with current technology, the use of CMCs in the hot section of the engine along with IMCs in the compressor is resulting in a 50% reduction in engine weight. This translates to an overall reduction in aircraft weight of nearly 40% for an aircraft with four engines, further contributing to lower initial costs, as well as lower operating costs.

The high-temperature composite materials required for these engines will have to operate satisfactorily from 5,000 to 16,000 hours at temperature. Interdiffusion, oxidation resistance, and creep, therefore, are major life-limiting problems that must be solved. Materials research also must include the study of failure modes and joining technology, and a mechanical and thermal-property database must be established. In addition, new, more precise design methods will be needed to address both the application of brittle composite materials and the integration of intricate cooling schemes for a wide range of material thermal conductivities. And finally, low-cost manufacture of the new materials and advanced components will require development of new fabrication processes.

Analytical modeling is being used to investigate the structural behavior of these advanced materials in six distinct areas: micro mechanics, deformation and damage, fatigue, fracture, trade-off studies, and loads definition. In the trade-off studies, coefficient of thermal expansion (CTE) mismatch, compliant layers, and fiber shape/size effects are being investigated using existing analytical tools to develop a physical understanding of advanced-composite development.

The emphasis in the area of loads definition is to develop and verify models to predict the aerodynamic and thermodynamic loads on a composite turbine blade. This is being accomplished by integrating existing aerodynamic, heat-transfer, and structural codes to predict blade response. The results are then calibrated and verified with simplified experiments that also are being defined and conducted under this task.

The results of analysis and experimental verification to date demonstrate the capability to simulate the high thermal gradients associated with engine operating conditions. In the future, this type of analysis will permit evaluation of an advanced-composite material`s performance in a simulated engine component.

Polymer-matrix composites (PMCs) are the lightest of the three types of composite materials under study in the HITEMP program. Recent applications of PMCs in aircraft propulsion systems, such as General Electric`s F-404 engine, have resulted in substantial reductions in both engine weight and manufacturing costs. Unfortunately, the low thermal-oxidation stability of PMCs severely limits the extent of their application. Commercially available state-of-the-art high-temperature PMCs, such as graphite fiber/PMR-15 and graphite fiber/PMR-11-55, are capable of withstanding thousands of hours of use at temperatures between 290 and 345°C).

To realize the full advantages of PMCs in aircraft-propulsion systems, however, new composite materials must be developed with enhanced thermal-oxidative stability permitting their use at temperatures to 425°C. Research on high-temperature PMCs under HITEMP is aimed at achieving this goal. Ongoing work includes:

   1. Study of the effects of resin/fiber interactions on composite stability and high-temperature performance
   2. Development of innovative processing techniques
   3. Exploration of oxidation-resistant coatings
   4. Synthesis of new polymers having good processability and significantly improved thermal-oxidative stability

Graphite-reinforced composites prepared with one of the new high-temperature polymers, V-CAP, undergo weight losses only about 60% those of comparable PMR-II-base composites after exposure in air at 370°C for 500 hours. An elevated-temperature nitrogen-postcure technique has been developed, which substantially improves the high-temperature (370°C) flexural strength of graphite-reinforced PMR-15 laminates. Application of this postcure method to V-CAP laminates enhances both the high-temperature mechanical properties and thermal-oxidative stability. Thus, the combined use of a higher stability matrix with improved processing yields a PMC with a useful lifetime in air at 370°C double that of a PMR-II-50 composite one of the best high-temperature PMCs currently available.

Continued improvements in the stability of polymer matrices coupled with improvements in polymer/fiber interfaces, composite processing, and oxidation-resistant coatings will yield PMCs for use at temperatures to 425°C.

Intermetallic-matrix composites. Several major problems limit the development of inter-metallic-matrix composites (IMCs), including chemical incompatibility and CTE mismatch between potential reinforcing fibers and matrix materials, poor low-temperature ductility, and marginal high-temperature oxidation resistance of intermetallic materials. Composite fabrication and joining processes that do not result in excessive fiber/matrix reaction or matrix contamination is an additional need.

The initial phase of the IMC program involves investigating available fiber compositions (SiC and Al2O3) in aluminides of iron, titanium, nickel, and niobium. These aluminides are Ti3Al and FeAl for applications to 1000°C and NiAl and Nb-alloy/aluminides for higher temperature applications. Alloying studies of these materials are aimed at increasing toughness, ductility, and oxidation resistance, and promoting longtime stability with the candidate fiber materials. Candidate matrices will be evaluated using tensile, compression, fatigue, creep, and oxidation tests. Measurement of appropriate thermal and physical properties is another planned task.

Powder-cloth fabrication processes have been developed to produce IMC materials, and alternative processing procedures, such as thermal spraying, are being studied. Encouraging results have been obtained on SiC-reinforced TiAl3 + Nb material, based on tensile, thermal-cycle, and strain-controlled fatigue studies for temperatures to 815°C.

The properties of first-generation SiC/Ti-24Al-11Nb composites compare favorably with those of current nickel-base super alloys on a strength/density basis. However, the SiC fiber is too reactive with the matrix material above 815°C, and also with the other candidate matrix materials. Therefore, researchers are focusing on using Al2O3 as the reinforcing fiber for these materials. There is a need for new fibers, however, and new compositions and fiber-processing techniques, such as the laser floating-zone process, have been identified. A project has been initiated to produce experimental quantities of fiber material.

Fiber coatings also are being investigated to function as diffusion barriers to limit fiber/matrix reaction and as compliant layers to lower stresses generated by CTE mismatch between the fiber and matrix. The oxidation resistance of FeAl is adequate for its intended use temperature and the time/temperature oxidation limits have been established for NiAl. Optimized fiber materials coupled with a better understanding of IMC behavior should result in future materials superior to those currently used for aerospace applications.

Ceramic-matrix composites. To meet HITEMP goals, CMC research is aimed at developing the basic and applied technologies needed to fabricate structurally reliable ceramic composites reinforced with long or continuous ceramic fibers. Like monolithic ceramics, these fiber-reinforced ceramics (FRCs) have lower densities, better oxidation resistance, and potential to operate at significantly higher temperatures than super alloys. However, unlike monolithic ceramics, FRCs display metal-like deformation behavior, noncatastrophic failure, and strength properties that is insensitive to processing- and service-generated flaws.

Recent investigations of a NASA-developed SiC/reaction-bonded silicon nitride (RBSN) composite system show that Si-based composite microstructures can be produced that are strong and tough for short times to temperatures.

Fiber development is critical since the development of advanced materials such as high-temperature composites is highly dependent on the availability of high-temperature fibers. If such advanced materials are going to be available for material-critical applications in future civil-transport engines, new fibers must be developed.

The wide range of fiber characteristics needed would require the development of more than one type of fiber. Fibers must have different properties, depending on the composite matrix, as well as the composite end use. In general, a candidate fiber should have low density, high strength, high stiffness, a CTE matching the matrix, chemical compatibility with the matrix, environmental stability, and appropriate fiber diameter.

The selection of appropriate fiber diameter also depends on the composite matrix. A large-diameter fiber (75 to 150 μm) is required for MMCs / IMCs to maximize fracture toughness. Small-diameter fibers ≤ 25 μm are required for CMCs to keep the critical flaw size for these brittle materials as small as possible. The environmental stability of the fiber also is a major factor; fibers must be able to withstand the high-temperature oxidation/hot-corrosion environment of the gas-turbine engine. This requirement emphasizes the need for the development of suitable fiber coatings, in conjunction with the development of the fibers themselves.

Fiber-research efforts begun under HITEMP include fiber fabrication by chemical vapor deposition, physical vapor deposition, polymeric precursors, and laser float-zone methods. Laboratory processes for fiber fabrication, however, are only the first steps toward the development of new high-temperature fibers. It is equally important to consider the scale-up required to produce the quantities of fiber needed for actual composite parts. A great deal of manpower and money is still required to scale-up from the small-size batches of fibers produced in the research laboratory to the vast quantities of fiber that will be needed in the future.

Formability or workability is generally defined as the amount of deformation that can be given to a specimen without fracture or necking in a given process. Workability is not an intrinsic material property; it depends on design variables:

    * process variables - stress, strain, strain rate, temperature, lubrication, etc., and
    * material variables - size, shape, and amount of second-phase particles, grain size, etc.

Therefore, for a given shape, workability is a function of material and process variables and can be expressed as

Workability = ƒ1 (material) -ƒ2 (process)

where ƒ1 is a measure of the ductility of the material under processing conditions, represented by forming limit criteria developed for various processes. Forming limit criteria based on limiting strains are of practical applicability because strains, as opposed to stresses, are easy to visualize and analyze in workability studies. The ƒ2 function, on the other hand, is given by stress, strain, strain-rate, and temperature histories at the potential failure sites of the work piece.

A complete workability analysis involves:

    * establishment of forming limit criteria (ƒ1) as a function of strain rate and temperature;
    * determination of stress, strain, strain-rate, and temperature histories (ƒ2) at potential failure sites; and
    * comparison of the results of flow analysis (ƒ2) with the forming limit criteria (ƒ1).

This comparison reveals the margin of safety for the deformation processing of a defect-free product. When a negative margin exists, it assists in deciding on the necessary changes in material or process variables, or both.

Free surfaces are the most commonly observed fracture sites in bulk deformation processes. In most cases, free-surface fractures determine the limits of deformation that can be imparted to the deforming material. Such fractures occur at the free surfaces of the specimen during processing, for example, edge cracking in rolling, surface cracking in bending, heading, open-die forging, or surface cracking before contact is achieved between the preform and the die walls in an impression-die forging.

There is developed a fracture criterion, based on limiting strains-to-fracture, for the prediction and prevention of surface cracks in bulk deformation processes. Local strains calculated from measurements, at fracture, of grid markings on the free surfaces of cylinders upset under different friction conditions and with different height-to-diameter ratios, are plotted. Fracture strains obtained from bend tests, measured by grid markings on convex surfaces of bend specimens, fall onto the extension of the fracture line determined by compression tests. Thus, bend tests are complimentary to compression tests, and are particularly useful when compression testing is not feasible.

The material function ƒ1, has not been studied systematically. Recently, has examined the effect of size, shape, and volume fraction of second-phase particles on the bulk formability of American Iron and Steel Institute (AISI) 1040, 1060, and 1090 carbon steels. The present study evaluates the workability of three heat-treatable aluminum alloys as influenced by aging and accompanying structural changes.

Three heat-treatable aluminum alloys (2014, 2024, and 7075) were received in the form of 12.7 mm diameter rods. The 2014 (≈100HB) and 2024 (≈130HB) aluminum alloys were in T4 condition. The 7075 alloy was received in T6 condition (≈150 HB).

Chemical compositions of aluminum alloys:

    * Aluminum alloy 2014: Cu-4.7%; Mg-0.5%; Mn-0.7%; Si-0.6%; Fe-0.3%
    * Aluminum alloy 2024: Cu-4.3%; Mg-1.5%; Mn-0.7%; Si-0.2%; Fe-0.3%
    * Aluminum alloy 7075: Cu-1.6%; Mg-2.5%; Mn-0.2%; Si-0.2%; Fe-0.3%; Zn-5.6%

The degree of banding in the 2014 aluminum alloy was more severe, and the elongated grains were larger near the surface. These large grains at the surface layers caused surface wrinkling during upset testing of the alloy, necessitating machining off a layer of 0.7 mm thickness from the surface in order to bring the wrinkling to an acceptable level. In the 2024 and 7075 aluminum alloys, surface wrinkling was minimal; the original surfaces were preserved during testing.

In order to study the effect of aging on workability, the alloys were solution treated (470-500°C) and aged to four different levels: naturally aged, peak-aged, over-aged, and highly over-aged. Solution treatments were carried out in a tube furnace in argon for the 2014 alloy, and in nitrogen for the 2024 and 7075 alloys. This was followed by quenching in an ice-water mixture.

Tests for the naturally aged condition were carried out after aging the specimens at room temperature (25°C) for one week. An oil bath was used for artificial aging. No recrystallization was detected after solution or aging treatments.

The 2024 aluminum alloy shows greater workability than the 2014 and 7075 alloys in all conditions. While the workability index in the 7075 alloy is improved ≈50 percent by over-aging, improvement in workability levels in the 2000 series is more pronounced. In these alloys, the workability index in the highly over-aged condition is approximately three times that of the naturally aged condition.

Upset test specimens of the 7075 alloy revealed exclusively 45-deg cracks in all conditions. The 2014 aluminum alloy specimens also showed 45-deg cracks, except in the highly overaged condition, where cracks in ≈20 percent of the specimens were vertical (also known as normal). These vertical cracks were randomly located on the fracture line.

In the 2024 alloy, both vertical and 45-deg cracks were observed in all conditions; in general, specimens with low aspect ratios tested under high friction conditions gave vertical cracks. The percentage of specimens containing vertical cracks increased with increased aging time. In the naturally aged condition, only ≈15 percent of the specimens had vertical cracks. This type of crack was seen in ≈20 percent in peak-hardness specimens and ≈50 percent of those in the over-aged conditions. In the highly over-aged condition, though, ≈90 percent of the specimens exhibited vertical cracks. It is clear that 45-deg cracking is the predominant fracture mode at low workability levels in this alloy.

The 45-deg cracks did not penetrate the cross-section of the specimens in the 2000 series alloys. Cracks in the 7075 alloy, however, generally traversed the cross-section of the specimen; there was no indication that cracking started in the center of the specimens.

In the three alloys, the poorest bulk workability was obtained in the naturally aged and peak-aged conditions, where the precipitated particles were small and sharable. Localization of shear in these conditions in heat-treatable aluminum alloys is well documented. Localization of shear and accompanying voiding in the 7000 series aluminum alloys has been studied and co-workers and Leroy and Embury.

Chung and co-workers observed the occurrence of localized shear failure in 7075-T4 aluminum alloy before the onset of necking and concluded that either deformation softening or negative strain-rate sensitivity was necessary for localization to occur. The degree of localization in overaged conditions, however, should be lower than that in the naturally aged condition, as evidenced by the small improvement in the workability level.

Results of tension and compression tests on 2014 alloy indicate that lack of shear localization in tension is not a guarantee that this phenomenon will be prevented in compression. In the compression of overaged specimens of this alloy, only 45-deg cracks are observed, and the persistence of localized shear failure is probable. It appears that the 2024 alloy is least affected by shear localization among the three alloys, as evidenced by the high degree of workability and occurrence of vertical cracks in all conditions.

This article covers various studies of thermo mechanical processing of AlMgZn alloys with particular emphasis on mechanical properties such as strength, ductility, fatigue and fracture toughness. The question as to whether thermomechanical processing improves stress-corrosion life of these alloys is critically examined. A systematic approach to a better understanding of the various stages of processing can be obtained through a correlation of microstructure and properties.

A critical problem encountered in the use of highest strength aluminum alloys (Al-Zn-Mg type) rests with an increased susceptibility to stress corrosion cracking as maximum strength levels are reached by conventional aging. Owing to the insidious nature of stress-corrosion failure, a compromise is made in the strength level at which the alloy is used, resulting in severe weight penalties, particularly in airframe applications.

In the past, variation of alloy-chemistry, solidification parameters, processing techniques and heat treatment have resulted in varying degrees of success toward product development in terms of strength and structural reliability. Of these, thermomechanical processing appears to be a most promising to the solution of this problem.

This article will only review the literature of thermo mechanical processing (TMP) of Al-Zn-Mg type alloy with particular emphasis on mechanical properties such as strength, ductility, fatigue and fracture toughness A systematic approach to a better understanding of the various stages of processing can be obtained through a correlation of microstructure and properties.

It is well known that there are strong interactions between the solute atoms and defects in aluminum alloys that result in structural instabilities variation solute profi1es and changes in solute diffusion rates. Investigations into repeated yielding in 7075 and other Al alloys by D’Antonio and others have shown that mobile dislocations are pinned by enhanced solute diffusion during tensile straining giving rise to the Portevin-LeChatelier (P-L) effect. This phenomenon appears in aluminum alloys over a range of temperatures above and below room temperature.

The amount of strain required for enhanced diffusion is a function of the temperature at which the straining is effected, and the strain-rate. Around room temperature, the diffusion coefficient of solute atoms in undeformed alloys is too low to result in their migration to dislocations.

If the diffusion of the solute atoms in a substitutional alloy is control1ed by a vacancy mechanism, then the diffusion coefficient, D, can be written as

D ≈ a2υZ Cv exp (-Em/kT)

where a is the interatomic distance, u is the average vibrational frequency of the atom, Z is the coordination number, Cv is the vacancy concentration, Em is the effective vacancy migration energy and k and T are the Boltzman constant and absolute temperature respectively.

To account for the enhanced diffusion, it has been proposed, that the diffusion coefficient is increased by the creation of non-equilibrium number of vacancies, Cv, due to plastic deformation, Cv can be related to the plastic strain Eo by the semiempirical relation

Cv = KEom

where K and m are material constants.

Plastic deformation not only increases the mobility of solute atoms in aluminum alloys but can cause clustering at dislocations. These clusters can as nucleation sites for subsequent strengthening precipitates. This phenomenon can be caused to occur at temperatures substantially below conventional aging temperatures and at which vacancy migration rates are low. This allows for wide variations in thermomechanical treatments and subsequent precipitate distribution and morphology.

Previous attempts to strengthen Al-Mg-Zn alloys thermomechanically have involved plastic deformation of solution-treated alloys prior to, during and subsequent to the aging process below, at and above room temperature.

McEvily have shown that 50% deformation at room temperature prior to aging produced slightly increased strength and greatly retarded crack growth rate during stress-corrosion. Morris found that large prestrains at room temperature produced instabilities that enhanced strengthening on subsequent aging. Post-aging deformation, i.e., cold working after stable precipitates have formed, also leads to strengthening.

DiRusso and others have found enhancement in both strength and ductility by deforming at an intermediate stage prior to attaining maximum strength in the aging process. Sommer tried to plastically deform in the peak hardened condition (T6) at an elevated temperature (above G.P. zone solvus temperature) and then overage to bring the strength level to T6 condition. The application of this TMP resulted in a material with the strength and fracture toughness of conventional T6 temper, possessing stress-corrosion cracking resistance equivalent to the overaged T73 condition.

The work on the effect of high temperature deformation in partially aged and in peak hardened condition toward property enhancement is quite extensive. However, none of these investigations have assessed in detail the effect of room temperature deformation on microstructure and properties in the underaged condition. These have led to a thermomechanical treatment yielding a product with superior physical and mechanical properties.

Osterman reported an increase of 25% in the 107 cycles-to-fai1ure stress level in 7075 Al. The above increase, over that of the conventionally aged alloy (T65l condition) was attained by thermomechanical treatment, where by the alloy is deformed at room temperature in a partially aged condition, and then post-deformation aged.

A great deal of controversy presently exists concerning the role of particular microstructural features on the stress-corrosion susceptibility of high strength alloy of 7075 type. Studies on stress-corrosion attack in Al-Zn-Mg type alloys have focused on three principal microstructural features:
a) matrix precipitate structure,
b) grain-boundary precipitate structure and
c) precipitate-free-zone (PFZ) which forms adjacent to the grain boundaries.

Work of Kent supports the Unwin and Nicholson contention that grain-boundary precipitate structure is of overriding importance to stress-corrosion attack. Other investigators contend that the nature of matrix precipitate is of importance to the deformation mode and thereby susceptibility. The planar slip mode is associated with susceptibility whereas materials deforming with a dislocation tangle structure is non-susceptible.

The importance of precipitate-free-zone has also been a point of controversy. The results of early work indicated that PFZ enhanced stress-corrosion susceptibility. More recent studies show that wide zones are preferred since they are weaker and undergo preferential deformation. Others contend that PFZ is not significant to stress-corrosion attack. All of the above three important features can be effected through thermomechanical treatments and hence have been used to improve stress-corrosion resistance.

Conclusion
A systematic study was undertaken to increase the yield strength and the resistance to stress-corrosion-cracking of 7075 A1 alloy beyond those of the T6 condition by thermomechanical processing. It has been found that yield strength over 85,000 psi which an accompanying elongation of about 11% and stress-corrosion-cracking resistance equiva1ent or superior to T73 condition can be achieved.

Thermomechanicat treatment involved 10% deformation of partially aged (3 days at room temperature and 5 hours at 120°C after solution treatment and water quench) samples and an aging treatment for 10 hours at 120°C. Samples which had been thermomechanically processed using the above sequence contained no precipitate-free-zones, precipitates of a size comparable to those in overaged condition and a stable and uniform dislocation substructure. Fracture toughness of the TMT material appears to be better than T6 but not as good as T73 condition.