KATALIN BALAZSI and CSABA BALAZSI

ABSTRACT

Biomaterials used for implant should posses some important properties in order to long-term usage in the body without rejection. One of the most important property is biocompatibility. These materials are used in different parts of the human body as artificial valves in the heart, stensts in blood vessels, replacement implant in shoul­ders, knees, hips and orodental structures. Materials used as different biomaterials should be made with certain properties as excellent biocompatibility, superior cor­rosion resistance in body environment, excellent combination of high strength and low modulus, high ductility and be without toxicity.

The creation of nanocomposites of ceramic materials with particle size few ten nanometers can significantly improve the bioactivity of the implant and enhance the osteoblast adhesion. One of the most used biomaterial is hydroxyapatite. The major inorganic constituent of bones and teeth is calcium phosphate, whose composition is similar to that of synthetic hydroxyapatite (HAp; Ca10(PO4)6OH)2. This similarity provides HAp based materials excellent bioactivity like bone bonding capability, osteoconductivity, and biocompatibility.

On the other hand, titanium (Ti) is most commonly used as orthopedic implant materials or bone substitute materials. Ti has good biocompatibility and sufficient mechanical properties for medical applications. One negative property of Ti is a low abrasion resistance and minute Ti abrasion powders may cause inflammatory reac­tions. Biomaterials must be chemically inert, stable and mechanically strong enough to wish stand the repeated forces a lifetime. From this point of view, TiC is a very stable phase in comparison to pure Ti or Ti alloys. Titanium carbide (TiC) is useful material for biomedical instruments because has a range of desirable properties. In this work, the combination of excellent bioactive hydroxyapatite with very stable and mechanically strong TiC has been studied. The nanostructured hydroxyapatite has been prepared by high efficient milling starting from biogenic eggshells. TiC

thin films were deposited by dc magnetron sputtering in argon atmosphere at differ­ent deposition temperatures. Spin coating was applied to obtain HAp decorated TiC films. Structural, mechanical and biological properties of HAp, Polymer-HAp and TiC-HAp coatings are being presented in this study.

2.1 INTRODUCTION

Biomaterials used for implant should posses some important properties in order to long-term usage in the body without rejection. One of most important properties is the biocompatibility. The biomaterial is “any substance, synthetic or natural in ori­gin, which can be used for any period of time, as a whole a part of a system which treats, augments or replaces any tissue, organ or function of the body.”1 Biomaterials are used in different parts of the human body as artificial valves in the heart, stents in blood vessels, replacement implant in shoulders, knees, hips and orodental structures.24 Materials used as different biomaterials should be made with certain properties. The materials used for orthopedic in plants should possess excellent biocompatibility, superior corrosion resistance in body environment, excellent combination of high strength and low modulus, high ductility and be without toxicity 5.

The materials currently used for implants include hydroxyapatite, 316L stain­less steel, cobalt-chromium alloys and pure titanium or its alloys. Elements such as Ni, Cr and Co are found to be released from the stainless steel and cobalt chro­mium alloys due to the corrosion in the body environment 6. The toxic effects of metals, Ni, Co and Cr released from prosthetic implants have been reviewed by Wapner7. Skin related diseases such as dermatitis due to Ni toxicity have been reported and numerous animal studies have shown carcinogenicity due to the pres­ence of Co 8.

The success of a biomaterial or an implant is highly dependent on three major factors; (i) the mechanical, tribological and chemical properties of the biomaterial, (ii) biocompatibility of the implant and (iii) the health conditions of the recipient and competency of the surgeon 9.

The biomaterials are grouped according to use in body. The situation is similar in the case of tissue. The tissue is grouped into hard and soft tissues. Tooth or bone are examples of hard tissue. Cartilage and ligament sor skin are the examples of soft tissues. These two types of tissues have the different properties from the structural or mechanical view. Considering the structural or mechanical compatibility with tissues, metal sor ceramics are chosen for hard tissue applications and polymers for soft tissue applications. The different mechanical properties of both types of tissues are shown in Table. 2.1.10.

In this chapter, THA hydroxyapatite based biomaterials developed as hard and soft tissue replacement were studied. The structural, mechanical and biological properties of bioinerttic—bioactive hydroxyapatite, biogen hydroxyapatite prepared from eggshells and polymer—hydroxyapatite composites were characterized.

When PGVNC is used as an epoxy hardener, it is supposed that all the phenolic hydroxy groups are hard to react with epoxy groups because of the steric hindrance. So, the epoxy/hydroxy ratio (1/1.14, 1/1.76, 1/2.65, 1/3.97) and curing temperature (150, 170, 190 °C) were optimized for the curing system of SPE and PGVNC. Table 4.9 summarizes the tan 5 peak temperature measured by DMA and 5% weight loss temperature of SPE/PGVNC cured at various conditions. When the epoxy/hydroxy ratio was changed at the fixed curing temperature of 170 °C which is a standard curing temperature of epoxy resin, SPE/PGVNC(1/2.65) had the highest tan d peak temperature, although 5% weight loss temperature decreased a little with decreas­ing epoxy/hydroxy ratio. When the curing temperature was changed between 150 and 190 °C at the fixed epoxy/hydroxy ratio of 1/2.65, the cured resin at 190 °C showed the highest tan 5 peak temperature (148.1 °C) and 5% weight loss tempera­ture (319.2 °C). This result suggests that 6.0 of 16 hydroxy groups of PGVNC are reacted with epoxy groups of SPE. For example, this number corresponds to the sum of two hydroxyl groups of four guaiacyl groups and four sets of one hydroxyl group per one pyrogallol unit in the guaiacyl pyrogallol[4]arene. Also, when the curing temperature is 190 °C, the tan d peak temperature of SPE/PGVNC 1/2.65 was high­er than that of SPE/PGVNC(1/1.76). We did not investigate the curing temperature higher than 190 °C, considering the stability of wood flour which is subsequently added. When SPE was cured with PN at 190 °C, SPE/PN(1/1) had a higher tan 5 peak temperature and 5% weight loss temperature than that of SPE/PN 1/2.65. This result is reasonable, considering that all the hydroxy groups of PN can react with the epoxy groups of SPE in contrast to the case of SPE/PGVNC.

Figure 4.39 shows the temperature dependency of E and tan 8 for SPE/ PGVNC(1/2.65) and SPE/PN(1/1) cured at 190 °C. The tan 8 peak temperature of SPE/PGVNC(1/2.65) (148. 1 °C) was much higher than that of SPE/PN(1/1) (78.1 °C). Also, the E of SPE/PGVNC(1/2.65) was higher than that of SPE/PNQ/1) over the temperature range from 0 to 200 °C. The fact that SPE/PGVNC has high glass transition temperature and rigidity should be attributed to the pyrogallol[4] arene structure. Figure 4.40 shows the comparison of tensile properties of SPE/ PGVNC(1/2.65) and SPE/PN(1/1). The SPE/PGVNC showed a higher tensile modulus than SPE/PN(1/1) in agreement with the result of DMA. However, tensile strength and elongation at break for SPE/PGVNC(1/2.65) were lower than those of SPE/PN(1/1), indicating a more brittle character due to the rigid calixarene struc­ture.

109 —

105 —

— 10

200

Temperature (°С)

The moisture absorption of jute composite samples were also measured according to ASTM D 5708. This test was carried out to determine the amount of moisture absorbed and the performance of the material in humid environment. Four square shaped samples of 3.5 cm x 3.5 cm of 5% natural rubber latex jute composite were tested. The samples were kept in an oven at temperature 100°C for one hour. Then the weights of the samples were measured. After that the samples were kept in a desiccator containing saturated solution of NaCl. The relative humidity of saturated Sodium chloride is 74.87 ± 0.12 at room temperature of 35°C. The percentage of moisture absorption of samples after 24, 48, and 120 h were calculated according to Eq. (4) as used in the water absorption test.

where W1 is original weight of the sample and W2 is final weight of the sample after 24, 48 and 120 h.

The results of moisture absorption test are shown in Table 6.3. The percentage of the moisture absorption gradually increases and then it reaches a saturation point and an average of 10.05% of moisture is absorbed approximately after 120 h.

7.3.2.1 TOW BUCKLING

Tow buckles mainly appear during the forming of reinforcement 1 on faces and edg­es of the tetrahedron shape. The localization of the buckles seems to be influenced by the initial positioning of the fabric, and the size of the buckles probably depends on the tension state of the tows perpendicular to the ones passing by the triple point. As a consequence, two initial positioning of the fabric have been tested as shown in Fig. 7.9 in conjunction with different blank holder pressures.

FIGURE 7.9 Initial positioning of the woven fabric.

Figure 7.10 shows that in the case of the orientation 0°, buckles only appear on edge 1 and on the middle of face 3. No buckles are observed on faces 1 and 2.

Fig. 7.10. Reinforcement 1: Localization of the buckle zone for initial fabric orientation of 0°

As the bending of the tows perpendicular to the ones passing by the triple point is the mechanisms supposed to be at the origin of the buckling defect, measurements of the bending angles on each faces has been carried out. Results are presented in Table 7.1:

TABLE 7.1 Bending Angle of the Horizontal Tows Measured on the Buckle Zone

The bending angles of the tows exhibiting buckling on the 3 faces of the shape were also measured. The values are reported in Table 7.2.

For orientation 90° the measured bending angles are situated in the same range of values as the ones measured for the 3 faces for orientation 0°. As a consequence,

the bending of the tows (Fig. 7.6d) is not responsible for the changes of the buck­le zone location for the 2 tested orientations. The initial reinforcement orientation seems to be crucial. As a consequence, bending is not a sufficient criterion to predict the appearance of the buckles.

The reinforcement considered in this study is not balanced. The tows, used in the warp and the weft directions are similar. However, a space between the weft tows (about the width of a tow) is observed on the fabric whereas this space is not present between the warp tows. As buckles only appear on bending zones where the weft tows are vertical, (face 3 and edge 1 orientation 0° and face 1 and face 2 orientation 90°) one can conclude that the architecture of the reinforcement is a key parameter conditioning the appearance of the buckles. When the warp tows are vertical (with­out any space between them) the buckles do not appear even though the horizontal tows exhibit the same amount of bending. This suggests that the presence of the space between the weft tows is one of the parameter that controls the appearance of the buckles.

The process parameters may also play a role in the occurrence of the buckling defect. The tows showing the buckles are not tight, and the effect of increasing the blank holder pressure upon the occurrence of the tow-buckling defect has been per­formed. In a first extent, the bending angles are considered. Table 7.3 reports for the 0° orientation the values of the bending angles in the 3 faces of the tetrahedron for three uniform blank holder pressures (uniform pressure applied to the fabric around the shape).

For reinforcement 1, Table 7.3 shows that the bending angle measured on Face Cslowly decreases as a function of the increasing blank-holder pressure. This is probably due an increasing deformation of the tows passing by the triple point as these ones can drag in a larger extent to the top of the shape the perpendicular tows showing the buckles. The relative similar values observed on Faces A and B are probably due to measurement dispersion and to their relative inaccuracy. For rein­forcement 2, no real tendency can be extracted from the bending angles values. It has to be noted that measurement were also performed for orientation 90° for both reinforcements and that similar conclusions can be emitted. As a consequence, the change of the blank holder pressure does not influence much the bending angle and therefore the tow orientation on the Faces and it is not really possible to control it.

However, the size, or height of the buckles may be influenced by an increasing tension of the tows. For reinforcement 2, small size buckles observed in Edge 1 dis­appear when a uniform blank holder pressure of 2 bar is applied. For reinforcement 1, the buckles remain. Table 7.4 shows the heights of the buckles, measured on edge 1 and face 3 for the 0° orientation and on faces 1 and 2 for the 90° orientation for different increasing uniform blank-holder pressures.

Table 7.4 shows that a relative reduction of the buckle size is observed on edge 1 while increasing the blank holder pressure. This reduction is probably due to a high­er tension in the tows showing the buckles. In Face C, the size of the buckles can be considered as constant as the precision of the measurement is about ± 0.1 mm.

To locally increase the tension on the tows showing the buckles, differential blank holder pressures can be applied. The pressure of blank holders 1 and 6 was increased with the goal to raise the tension in the tows showing the buckles in edge 1 (opposite of Face C). The pressure in the other blank holders remains at 1 bar.

TABLE 7.5 Size of the Buckles as a Function of the Increasing Blank Holder 1 and 6 Pressures

Table 7.6 shows that the size of the buckles this time decreases in Face C by increasing the tension in blank holder 2 and 4 as the tension in the tows showing the buckles is raised. However, the size of the buckles in Edge 1 remains constant.

The observations performed on the two previous tests indicate that it may be difficult to decrease simultaneously the size of the buckles and therefore to stop their occurrence by only working with the blank holder pressure and this for our test configuration.

It has to be noticed that the blank-holder pressure was not raised above values of

2.5 bar as another defect (sliding of tows within the membrane) appears in this case.

The sugar palm fiber (SPF) was collected at Jempol, Negeri Sembilan in Malaysia. All of the fibers were grind and screened using Fritsch pulverisette mill to obtain 2 mm fiber size. For the extraction of sugar palm starch (SPS), firstly, the woody fibers and starch powder was obtained from the interior part of the trunk (see Fig. 9.13). Then, this mixture (woody fibers and starch powder) was carried out for washing process to obtain the starch. Then starch was kept in an open air for a mo­ment and dried in an air circulating oven at 120°C for 24 hrs.

FIGURE 9.13

The effect of the reinforcement agent on the notched Izod impact energy for the composite is also listed in the following table.

The average for five test specimens and their significant standard error is given for each property. Figures 12.7 and 12.8 graphically summarize the data listed in Table 12.3.

matrix. This behavior is observed for all matrix reinforcement combinations, al­though the rate in reduction of the tensile strength and Izod impact strength varied from case to case, depending on the reinforcement. Most of all plant fibers are hy­drophilic in nature with a moisture content enough high due to the presence of cel­lulose in cell structure. All these organic reinforcements generally have high aspect ratio, so, the efficiency of transmitting stress from matrix to these types of agents is quite poor. On the other hand, lignin increased the hard segments of composite the films, making the films less elastic and more brittle, which led to the impact strength decreasing. This explains why the impact strength decreased as the filler content reached 40% in weight. Poor interfacial bonding causes partially separated micro spaces between the filler and the matrix polymer, which obstructs stress propaga­tion, when tensile stress is applied, and induces decreased strength and increased brittleness but compatibilizing agent can solve partially this problem. This is what justifies the use of a coupling agent such as MAPP in all formulations. As the rein­forcement loading (50%) is higher, filler-filler agglomeration occurs and degree of weak interface regions between reinforcement particles and matrix become more and leads deterioration in tensile strength. As the degree of agglomeration increases, the filler- matrix interaction becomes poor, leading to the decrement in the tensile strength. Incorporation of both types of fillers (rice, wood and DDGS), generates a reduction in tensile strength of PP and its composites.

Figure. 12.7(b) shows that notched impact strength decreases in all the com­posites. This may be explained by the fact that the presence of reinforcing fillers ends within the body of the composite can cause crack initiation and subsequent failure. The reason is that the ends of reinforcing fillers act as notches and generate considerable stress concentrations, which could initiate micro cracks in the ductile PP matrix.

The impact test machine used in this study did not provide enough energy to break the neat PP because of the high flexibility of the PP matrix. By contrast, all specimens broke completely into two pieces. Introducing the reinforcement in the composites led to an increased stress concentration because of the poor bonding between the reinforcement (wood, rice or DDGS) and the polymer. As impact wave met different phases such as fiber, polymer, and voids in the cross machine direction, it would lose its energy as dissipation energy. Although crack propagation became difficult in the polymeric matrix reinforced with filler, the decrease in the impact energy observed was ascribed to fiber ends, at which micro cracks formed and fibers debonded from the matrix. These micro cracks were a potential point of composite fractures. Another reason for the decreased impact strength may have been the stiff­ening of polymer chains due to the bonding between the wood fibers and the matrix. For high-impact properties, in fact, a slightly weaker adhesion between the fiber and polymer is desirable, as it results in a higher degradation of impact energy and sup­ports the so-called fiber pullout56. In composites, the effect of the reinforcement is to increase the tendency to agglomerate, which generates a low interfacial adhesion

leading to the weakening of the interfacial regions. These agglomerates then act as sites for crack initiation. Poor interfacial bonding has been indicated in the literature as the major reason for the loss in strength and elastic modulus57.

Adding fillers also resulted in an increment of void content, which contributes to stress concentration, thus reducing strength. This behavior is consistent with what is observed in the impact tests that revealed a decrease in composites samples. The presence of numerous cavities is clearly visible in Fig. 12.4b (PWM) which has the lowest impact strength, this indicates that the level of interfacial bonding between the fibers and the matrix is weak and when stress is applied it causes the fibers to be pulled out from the matrix easily leaving behind gaping holes. These two properties are indicators of the plasticity of the material, and showed that the PP has a tendency for the occurrence of fracture with loading reinforcement. The flexural modulus of composites is influenced mainly by the adhesion between the matrix and disper­sion of reinforcing fillers inside it. The results for this mechanical property also supported the existence of a certain degree of miscibility in the composite plastics (Fig. 12.8).

From Fig. 12.8, all the compositions showed aflexural modulus higher than the pure PP. This increase in flexural properties was expected due to the improved ad­hesion between components in the blends. For some authors58, this is due to the restriction of the mobility and deformability of the matrix with the introduction of mechanical restraint. Many researchers59,60,61 have observed that the inclusion of wood fibers or lignocellulosic fibers into thermoplastics such as, polyethylene, or PP generally results in a decrease in tensile strength and elongation at break but an increase in Young’s modulus. This increase of flexural modulus can be attributed to the increase in volume fractions of high-modulus fibers in plastic composites62. When increasing the reinforcement, tensile and compression strengths constantly decreased. The presence of the wood or other reinforcement in the polymeric matrix augmented the polymer’s rigidity, increasing the value of the modulus in relation to the pure polymer. This phenomenon has also been reported by other research­ers who studied the effect of wood flour on mixtures of recycled polystyrene and polyethylene63,64.

and virgin polystyrene65. The Flexural modulus of composite with DDGS is sig­nificantly higher than all the other ones, may be, this is likely due to the oils being removed in the DDGS material. So, we can notice that the initial chemical treat­ments on DDGS certainly has a positive effect on the improvement of mechanical properties (FM and IS) that are higher than those of untreated wood. These results are in agreement with those of Julson et al61. The improved dispersion obtained from the composite with treated DDGS was also responsible for the highest flex­ural modulus. For none coupled composite (wood and PP only), the filler particles began to form aggregates. Direct physical bonds between filler particles are weak and, thus, easily broken during tensile loading, which explains the decrease in the flexural modulus (WPM). Compatibilizers can change the molecular morphology of the polymer chains near the fiber-polymer interphase. Yin et al.66 reported that the addition of coupling agent (MAPP) even at low levels (1-2%) increases the nucle — ation capacity of wood-fibers for polypropylene, and dramatically alters the crystal morphology of polypropylene around the fiber. When MAPP is added, surface crys­tallization dominates over bulk crystallization and a transcrystalline layer can be formed around the wood-fibers. Crystallites have much higher moduli as compared to the amorphous regions and can increase the modulus contribution of the polymer matrix to the composite modulus67. The flexural modulus of the composites can be correlated with the morphology of these ones. Composites whose surfaces are smoother and more homogeneous exhibit the greatest flexural modulus. The resul­tant increase in flexural modulus properties (Fig. 12.8) can be explained on the basis of improved wettability (compatibility) of the reinforcement fibers with the polymer matrix. The increased compatibility is obtained by reducing the polarity of the wood fiber surface nearer to the polymer matrix.

The mechanical results of this study show that loading of PP with these natural fibers leads to a decrease in tensile and impact strength of the pure polymer. On the other hand, the flexural modulus increases due to the higher stiffness of the fibers. The significant improvements in flexural properties of the blends composites made with MAPP and reinforcing fillers were further supported by SEM micrographs.

12.2 CONCLUSION

Wood fibers, rice husk and DDGS, which originate from renewable resources, are an interesting alternative to mineral fibers. All these samples with reinforcement exhibited a markedly heterogeneous and highly rough fracture surface with large voids, or cavities, around the filler particles due to the accumulation of stresses in the particle-matrix interface zone. This produced an adverse effect on mechanical properties such as tensile strength and impact resistance.

The SEM micrographs reveal that interfacial bonding between the treated filler and the matrix has significantly improved, suggesting that better dispersion of the filler into the matrix was achieved upon treatment of rice husk and DDGS. The ther­mal properties revealed the strong nucleation ability of the reinforcement flour and MAPP on PP crystallization. Crystallization of all the composites with the coupling agent MAPP only or with reinforcement began earlier compared to that of pure PP. This suggested that MAPP and organic reinforcement acted as nucleation agents and were responsible of the shift of crystallinity towards higher temperatures.

Tensile and impact strength exhibited a marked downward tendency as reinforce­ment was loaded. This is due to the weak interfacial adhesion and low compatibility between matrix and filler. The weak bonding between the hydrophilic lignocellu- losic agent and the hydrophobic matrix polymer obstructs the stress propagation and causes the decrease of these properties. These properties were not significantly affected by changing rice husk by DDGS. However, the flexural modulus for all composites with organic reinforcement is higher than the value for neat PP, as a consequence of the high modulus of cellulosic agent.

In summary, the use of organic reinforcements as fillers to polymer matrix com­posites proved to be a viable alternative. Reductions in tensile and impact strength properties reported with the addition of fillers may be tolerable to some applica­tions. Increment in flexural modulus is achieved in all cases. The development of alternatives for recycling rice husks and DDGS as reinforcements in polymer matrix composites is an important step to provide a good destination for these wastes and opens an opportunity to produce a new value added product. This development can permit reduction of production costs and less use of wood in composites based on polymer matrix, especially with the scarcity of wood across the world. This novel application of rice husk and DDGS for bio composites has significantly higher eco­nomic value than its traditional use as a feed stuff. The distillers’ dried grains with soluble (DDGS) from corn ethanol industry; the rice husk products show immense opportunities in engineering new green composites when integrated with thermo­plastics. The properties of PP composites can be adjusted by mixing different rein­forcements species for filler blend.

12.3 FUNDING

This research received no specific grant from any funding agency in the public, commercial or not-for-profit sectors.

• Melting Point: —115 °C

• Crystallinity: Low crystallinity (50-60% crystalline). Main chain contains many side chains of 2-4 carbon atoms leading to irregular packing and low crystallinity (amorphous).

• Strength: Not as strong as HDPE due to irregular packing of polymer chains.

• Transparency: Good transparency, since it is more amorphous (has noncrys­talline regions) than HDPE.

• Density: 0.91-0.94 g/cm3, lower density than HDPE.

• Chemical properties: Chemically inert. Insolvent at room temperature in most solvents. Good resistance to acids and alkalis. Exposure to light and oxygen results in loss of strength and loss of tear resistance.

• Tensile elongation at rupture (%): 906.

15.1.2.3 APPLICATIONS

Low-density polyethylene (LDPE) is used mainly in film applications for both pack­aging and nonpackaging applications (Fig. 15.3). Other markets include extrusion coatings, sheathing in cables and injection molding applications. LDPE is the oldest and most mature of the polyethylenes (PEs). It is characterized by its short and long chain branching, which gives it good clarity and processability although it does not have the strength properties of the other PEs.

FIGURE 15.3 LDPE application.

MFC is normally produced from highly purified wood fiber (WF) and plant fiber (PF) pulps by high pressure homogenization according to the procedures developed at ITT Rayonnier.6263 Pulp is produced by using a mixture of sodium hydroxide and sodium sulfide and thus so-called Kraft pulp (almost pure cellulose fibers) is obtained. Pulping with salts of sulfurous acid leads to cellulose named sulphite pulp (which contains more by-products in the cellulose fibers). MFC particles are con­sidered to comprise of several elementary fibrils. Each one of them consisting of 36 cellulose chains has a high aspect ratio or ~10-100 nm wide and 0.5-10 pm in length. MFCs are ~100% cellulose, and contain both amorphous and crystalline regions. In food and cosmetic industries, MFCs have been used as a thickening agent.64

TYLER CHUANG, ALBERT LIN, and CRAIG VIERRA

ABSTRACT

Spider silk has extraordinary mechanical properties, containing a unique balance of high-tensile strength and extensibility. Spider silks outperform several well-known manmade materials including high-tensile steel, Kevlar (body armor), and nylon. Modern spiders spin at least six to seven different fiber types that have distinct me­chanical properties. These fibers have been shown to be biocompatible, heat stable and environmentally green materials. Because of their unique mechanical features, scientists are pursuing spider silks as next generation biomaterials that can be used for a broad range of applications. Given the cannibalistic nature of spiders, com­bined with their venomous nature, farming spiders for silk becomes a dangerous and impractical method for obtaining large amounts of silk for industrial applications. This has prompted scientists to develop DNA methodologies and heterologous pro­tein expression systems that produce vast quantities of recombinant silk proteins. Some potential applications of silks spun from recombinant proteins include body armor, ropes and cords, sutures, tissue scaffolds, tires, and shoes. Here we cover the following topics: the diverse protein nature of spider silks and the compositions of threads, the method and steps for synthetic silk production, and the translation of fiber production from a small scale to a large-scale spinning platform. By drawing upon advances in spinning methodologies, we will describe the development of spider silk fiber composites using full-length silk proteins retrieved from the major ampullate silk-producing gland blended with other recombinant proteins expressed and purified from microorganisms.

2.4.1.1 PREPARATION OF TIC/A:C THIN FILMS

Modern methods of vacuum deposition provide great flexibility for manipulating ma­terial chemistry and structure, leading to films and coatings with special properties. TiC/A:C nanocomposite thin films have been prepared by DC magnetron sputter­ing (Fig. 2.15) on silicon (001) substrate with 300 nm thick oxidized silicon sub­layer. Films have been deposited at 200 °C in argon at 0.25 Pa. The input power of 99.999% purity carbon target (C) was kept constant at 150 W and the input power of the 99.995% purity titanium target (Ti) was changed between 15 and 50 W. The deposition rate was ~0.06 nm/s. The thickness of thin films was about 300 nm. The structural investigations have been performed on a Philips CM-20 Microscope using a 200 kV accelerating voltage. The elemental analysis of film composition was also performed in this microscope, which is equipped with a NORAN EDS (Energy Dis­persive Spectrometer), with an HP-Ge detector.