*Steelmaking*, Vol. 37, 2010, pp.313-319

Vol. 38, 1998, pp. 16-22

pp.435–440

http:/rsb.info.nih.gov/ij

Inspection System for Pseudo-particle Analysis. *Iron and Steel Institute of Japan*

Bearing Resources on lron Ore Sintering. *Iron and Steel Institute of Japan International*,

iron ore sintering process. *Resources, Conservation and Recycling*, Vol.55, 2011,

tomography to investigation of structure of sinter mixture granules. *Ironmaking &*

porosity characterisation with application of 3D X-ray tomography. *Ironmaking &*

Many applications on modern technologies try to combine uncommon properties that can't be attended by conventional metallic alloys, ceramics and polymeric materials. On this way, engineers of aerospace, for example, are ever looking for structural materials that have low densities, be stronger, with wear resistance and that have high resistance to corrosion (Gordo et al., 2000). These characteristics combine amazing properties for just one material. Normally, very dense materials are harder with high resistance to wear. Thus, composite materials are designed to provide this mixing of controlled properties.

The principle of combined action of different properties from two or more distinct materials characterizes the composite materials. Then, it's the central principle of composite materials that is enhancing one or more properties my mixing materials in various ways. Thus a very elementary example of a composite would be mud mixed with straw; still a very widely used material in the construction of houses. These innovative materials open up unlimited possibilities for modern material science and development.

Basically there are three kinds of composite materials, they are: polymer matrix composite (PMC), for example fiberglasses embedded in a thermoplastic resin of polyethylene; ceramic matrix composite (CMC) with carbon fibers reinforcing SiC matrix; and metallic matrix composite (MMC) to reinforce a matrix of steel with refractory ceramic hard particles.

This last one kind of composite material will be discussed in this chapter with focus on effect of the dispersion of nanosized carbides (NbC and TaC) in the sintered microstructure of the stainless steel 316L. Powder metallurgy technique will be applied as well as mechanical alloying and the effects promoted by hard particles in a metal matrix. A practical case study will be shown.

#### **2. Metal Matrix Composites (MMC)**

Metal matrix composites are an attractive end goal for many materials engineers as an alternative to the traditional pure ceramics or pure metal material. The goal is to create a

Effect of the Additives of Nanosized Nb and Ta Carbides

material. The following demands are generally applicable:

• Low density,

the matrix),

• Chemical compatibility, • Thermal stability, • High Young's modulus,

• Good processability, • Economic efficiency.

(Torralba et al., 2003).

lamination).

• High compression and tensile strength,

on Microstructure and Properties of Sintered Stainless Steel 235

For other applications different development objectives are given, which differ from those mentioned before. For example, in medical technology, mechanical properties, like extreme corrosion resistance and low degradation as well as biocompatibility are expected. Although increasing development activities have led to system solutions using metal composite materials, the use of especially innovative systems, particularly in the area of light metals, has not been realized. The reason for this is insufficient process stability and reliability, combined with production and processing problems and inadequate economic efficiency. Application areas, like traffic engineering, are very cost orientated and conservative and the industry is not willing to pay additional costs for the use of such materials (Kainer, 2006). Reinforcements for metal matrix composites have a manifold demand profile, which is determined by production and processing and by the matrix system of the composite

• Mechanical compatibility (a thermal expansion coefficient which is low but adapted to

On the other hand, the matrix has the role of distribute and transfer the tensile and to protect the surface of reinforcement. Reinforcements can be used as particles, whiskers or fibers. Unfortunately, high strength fibers and their processing methods are extremely expensive, and this limits their wide industrial application. Moreover, continuous fiber reinforced composites do not usually allow secondary forming process, which is used in the original shape in which they were manufactured. As a result of these limitations, over the last few years, new efforts on the research of non-continuous reinforced composites have been applied. Composites with non-continuous reinforcement do not have the same level of properties as continuously reinforced composites, but their cost is lower, their processing methods are more adaptable to conventional ones and their performance is acceptable

The selection of suitable matrix alloys is mainly determined by the intended application of the composite material. With the development of light metal composite materials that are mostly easy to process, conventional light metal alloys are applied as matrix materials. In the area of powder metallurgy special alloys can be applied due to the advantage of fast solidification during the powder production. Those systems are free from segregation problems that arise in conventional solidification. Also the application of systems with

Normally, the process to produce MMC involves at least two steps: consolidation and synthesis (i.e. introduction of the reinforcement into the matrix), following by a modeling operation. There are a lot of techniques to consolidate; MMC reinforced with discontinuous fibers are susceptible to modeling through mechanical forming (i.e. forging, extrusion and

oversaturated or metastable structures is possible (Kainer, 2006).

material that has nearly the same high strength as the ceramics, but also high fracture toughness due to the contribution from the metal, this has lead to the development of a new family of metallo-ceramic composites. Classical examples are the precipitation and nucleation of carbides in steel for obtaining good strength while still maintaining most of the toughness from the metal, or the fabrication of hard metals such as a matrix cobalt containing a large percentage of grains of hard tungsten or titanium carbides or nitrides, which during the powder processing grow together and form a continuous skeleton (Breval, 1995).

In principle, wrought irons and most conventional steels could be treated as MMCs since they have a metallic matrix which is reinforced with a dispersed phase that could be oxides, sulphides, nitrides, carbides, etc. Even if we take a definition through "microstructural design", many conventional engineering alloys, such as some steels involving dispersions created at a moving alpha-gamma interface, really ought to be included. Although, composite material is a definition very embracing, it is quite tendentious to think that a directional solidification of a eutectic microstructure is within the framework of an MMC definition (Ralph, 1997).

The reinforcement of metals can have many different objectives. The reinforcement of light metals opens up the possibility of application of these materials in areas where weight reduction has first priority. The precondition here is the improvement of the component properties. The development objectives for light metal composite materials are:

To summarize, an improvement in the weight specific properties can result, offering the possibilities of extending the application area, substitution of common materials and optimization of component properties. With functional materials there is another objective, the precondition of maintaining the appropriate function of the material. Objectives are for example:

For other applications different development objectives are given, which differ from those mentioned before. For example, in medical technology, mechanical properties, like extreme corrosion resistance and low degradation as well as biocompatibility are expected. Although increasing development activities have led to system solutions using metal composite materials, the use of especially innovative systems, particularly in the area of light metals, has not been realized. The reason for this is insufficient process stability and reliability, combined with production and processing problems and inadequate economic efficiency. Application areas, like traffic engineering, are very cost orientated and conservative and the industry is not willing to pay additional costs for the use of such materials (Kainer, 2006).

Reinforcements for metal matrix composites have a manifold demand profile, which is determined by production and processing and by the matrix system of the composite material. The following demands are generally applicable:

• Low density,

234 Sintering – Methods and Products

material that has nearly the same high strength as the ceramics, but also high fracture toughness due to the contribution from the metal, this has lead to the development of a new family of metallo-ceramic composites. Classical examples are the precipitation and nucleation of carbides in steel for obtaining good strength while still maintaining most of the toughness from the metal, or the fabrication of hard metals such as a matrix cobalt containing a large percentage of grains of hard tungsten or titanium carbides or nitrides, which during the powder processing grow together and form a continuous skeleton (Breval,

In principle, wrought irons and most conventional steels could be treated as MMCs since they have a metallic matrix which is reinforced with a dispersed phase that could be oxides, sulphides, nitrides, carbides, etc. Even if we take a definition through "microstructural design", many conventional engineering alloys, such as some steels involving dispersions created at a moving alpha-gamma interface, really ought to be included. Although, composite material is a definition very embracing, it is quite tendentious to think that a directional solidification of a eutectic microstructure is within the framework of an MMC

The reinforcement of metals can have many different objectives. The reinforcement of light metals opens up the possibility of application of these materials in areas where weight reduction has first priority. The precondition here is the improvement of the component

• Increase in yield strength and tensile strength at room temperature and above while

• Increase in creep resistance at higher temperatures compared to that of conventional

To summarize, an improvement in the weight specific properties can result, offering the possibilities of extending the application area, substitution of common materials and optimization of component properties. With functional materials there is another objective, the precondition of maintaining the appropriate function of the material. Objectives are for

• Increase in strength of conducting materials while maintaining the high conductivity,

• Improvement in low temperature creep resistance (reactionless materials),

• Production of layer composite materials for electronic components,

• Increase in operating time of spot welding electrodes by reduction of burn outs,

• Improvement of burnout behavior (switching contact), • Improvement of wear behavior (sliding contact),

• Production of ductile composite superconductors,

• Production of magnetic materials with special properties.

properties. The development objectives for light metal composite materials are:

maintaining the minimum ductility or rather toughness,

• Increase in fatigue strength, especially at higher temperatures,

• Improvement of thermal shock resistance, • Improvement of corrosion resistance,

• Increase in Young's modulus, • Reduction of thermal elongation.

1995).

definition (Ralph, 1997).

alloys,

example:

On the other hand, the matrix has the role of distribute and transfer the tensile and to protect the surface of reinforcement. Reinforcements can be used as particles, whiskers or fibers. Unfortunately, high strength fibers and their processing methods are extremely expensive, and this limits their wide industrial application. Moreover, continuous fiber reinforced composites do not usually allow secondary forming process, which is used in the original shape in which they were manufactured. As a result of these limitations, over the last few years, new efforts on the research of non-continuous reinforced composites have been applied. Composites with non-continuous reinforcement do not have the same level of properties as continuously reinforced composites, but their cost is lower, their processing methods are more adaptable to conventional ones and their performance is acceptable (Torralba et al., 2003).

The selection of suitable matrix alloys is mainly determined by the intended application of the composite material. With the development of light metal composite materials that are mostly easy to process, conventional light metal alloys are applied as matrix materials. In the area of powder metallurgy special alloys can be applied due to the advantage of fast solidification during the powder production. Those systems are free from segregation problems that arise in conventional solidification. Also the application of systems with oversaturated or metastable structures is possible (Kainer, 2006).

Normally, the process to produce MMC involves at least two steps: consolidation and synthesis (i.e. introduction of the reinforcement into the matrix), following by a modeling operation. There are a lot of techniques to consolidate; MMC reinforced with discontinuous fibers are susceptible to modeling through mechanical forming (i.e. forging, extrusion and lamination).

Effect of the Additives of Nanosized Nb and Ta Carbides

(Suryanarayana, 2001).

promote an effect of hardening.

**5. Mechanism of reinforcement**

Clyner & Withers, 1993 and McDaniels et al., 1965).

on Microstructure and Properties of Sintered Stainless Steel 237

intermetallics and mechanochemical synthesis of materials have been critically reviewed after discussing the process and process variables involved in MA (Suryanarayana, 2001). The process of MA consists of loading the powder mix and the grinding medium (generally hardened steel or tungsten carbide balls) in a container sealed usually under a protective atmosphere (to avoid/minimize oxidation and nitridation during milling) and milling for the desired length of time. There are some parameters that must be controlled in the milling process, such as ball to powder ratio, length of milling time, atmosphere and ball sizes

Nowadays, mechanical alloying has been successfully applied to prepare Ni-based, Febased, and Al-based oxide dispersion strengthened (ODS) alloys and carbide dispersion strengthened (CDS) alloys. During preparation of these materials, the strengthened phase is directly put into the matrix powders as a raw material. After a long time of ball milling, mechanical alloying can be achieved to some extent, but it is unlikely to produce the optimum combination between the strengthened phase, and the matrix. Element powders which can form oxide or carbide are mixed with the matrix powders and then subjected to ball milling, the supersaturated solid solution will form first, from which the dispersion strengthened alloy can be fabricated by a precipitation of second phase during heat treatment, or by mechanical-chemical reaction. This is a new method to produce CDS and ODS alloys by mechanical alloying, which can be used to optimize the strengthened action of the second phase to the maximum (Wang et al., 2001). During mechanical alloying, the mechanical energy transfers from the balls to the alloy powders, which results in the severe plastic deformation and the introduction of dislocations, lattice distortion and vacancies. When hard particles are incorporated into the matrix they

The characteristics of metal matrix composite materials are determined by their microstructure and internal interfaces, which are affected by their production and thermal mechanical prehistory. The microstructure covers the structure of the matrix and the reinforced phase. The chemical composition, grain and/or sub-grain size, texture, precipitation behavior and lattice defects are of importance to the matrix. The second phase is characterized by its volume percentage, its kind, size, distribution and orientation. Locally varying internal tension - due to the different thermal expansion behavior of the two phases - is an additional influencing factor. With knowledge of the characteristics of the components, the volume percentages, the distribution and orientation, it might be possible to estimate the characteristics of metallic composite materials. The approximations usually proceed from ideal conditions, i.e. optimal boundary surface formation, ideal distribution (very small number of contacts of the reinforcements among themselves) and no influence of the component on the matrix (comparable structures and precipitation behavior). However, in reality a strong interaction arises between the components involved, so that these models can only indicate the potential of a material. The different micro-, macro- and meso-scaled models proceed from different conditions and are differently developed (Kainer, 2006 as cited by

#### **3. Powder Metallurgy (P/M) used to obtain composite materials**

Between lots of ways to produce composite materials, powder metallurgy (P/M) is an easy way to manufacture composite materials with metallic matrix and offer some advantages compared with ingot metallurgy or diffusion welding. The main advantage is the low manufacture temperature that avoids strong interfacial reaction, minimizing the undesired reactions between the matrix and the reinforcement. In other cases, P/M allows production of some materials which cannot be obtained by any other alternative route (for example, SiC reinforcing Ti alloys). Composites that use particles or whiskers as reinforcement can be obtained easier by P/M than by other routes. Another advantage of P/M is the homogeneous distribution of reinforcements. This uniformity not only improves the structural properties but also the reproducibility level in properties (Torralba et al., 2003).

In short, powder metallurgy is the process of blending fine powdered materials, pressing them into a desired shape (compacting), and then heating the compressed material in a controlled atmosphere to bond the material (sintering). Sintering is a term applied to bond the material and it describes the process in which aggregated or compacted powders acquire a solid structure, coherent through internal displacements of mass by diffusion process of atoms activated by temperature (Gomes, 1995). P/M process generally consists of four basic steps: (1) powder manufacture, (2) powder blending, (3) compacting and (4) sintering. Compacting is generally performed at room temperature, and the elevated-temperature process of sintering is usually conducted at atmospheric pressure. Hot Isostatic Pressure (HIP) is another way to press and to sinter at the same time. This technique avoids some troubles in the pressing step. Optional secondary processing often follows to obtain special properties or enhanced precision such as thermal treatments. Each step has specific parameters that must be studied and controlled to reach the performance desired.

Among some techniques to manufacture powders, such as liquid metal atomization, chemical and thermal decomposition reactions, electrolyte deposition, sol-gel process and milling, this last one will be shortly discussed in the next section.

#### **4. Mechanical alloying and milling**

Two kinds of millingcould be distinguished - conventional low energy milling ( e.g. in a ball mill) and high energy milling. In the last case, the most used devices are planetary ball mill, SPEX ball mill and attritor ball mill. High energy milling (HEM) and mechanical alloying (MA) are terms utilized to name the synthesis of alloys and compounds by means of intensive milling that causes mixing at an intimate level and deformation of the materials microstructure (Costa et al. as cited by Benjamin, 1976 and Suryanarayana, 2001). MA is a solid-state powder processing technique involving repeated welding, fracturing, and rewelding of powder particles in a high-energy ball mill. Originally developed to produce oxide-dispersion strengthened (ODS) nickel- and iron-base superalloys for applications in the aerospace industry, MA has now been shown to be capable of synthesizing a variety of equilibrium and non-equilibrium alloy phases starting from blended elemental or prealloyed powders. The non-equilibrium phases synthesized include supersaturated solid solutions, metastable crystalline and quasicrystalline phases, nanostructures, and amorphous alloys. Recent advances in these areas and also on disordering of ordered

Between lots of ways to produce composite materials, powder metallurgy (P/M) is an easy way to manufacture composite materials with metallic matrix and offer some advantages compared with ingot metallurgy or diffusion welding. The main advantage is the low manufacture temperature that avoids strong interfacial reaction, minimizing the undesired reactions between the matrix and the reinforcement. In other cases, P/M allows production of some materials which cannot be obtained by any other alternative route (for example, SiC reinforcing Ti alloys). Composites that use particles or whiskers as reinforcement can be obtained easier by P/M than by other routes. Another advantage of P/M is the homogeneous distribution of reinforcements. This uniformity not only improves the structural properties but also the reproducibility level in properties (Torralba et al., 2003).

In short, powder metallurgy is the process of blending fine powdered materials, pressing them into a desired shape (compacting), and then heating the compressed material in a controlled atmosphere to bond the material (sintering). Sintering is a term applied to bond the material and it describes the process in which aggregated or compacted powders acquire a solid structure, coherent through internal displacements of mass by diffusion process of atoms activated by temperature (Gomes, 1995). P/M process generally consists of four basic steps: (1) powder manufacture, (2) powder blending, (3) compacting and (4) sintering. Compacting is generally performed at room temperature, and the elevated-temperature process of sintering is usually conducted at atmospheric pressure. Hot Isostatic Pressure (HIP) is another way to press and to sinter at the same time. This technique avoids some troubles in the pressing step. Optional secondary processing often follows to obtain special properties or enhanced precision such as thermal treatments. Each step has specific

parameters that must be studied and controlled to reach the performance desired.

milling, this last one will be shortly discussed in the next section.

**4. Mechanical alloying and milling**

Among some techniques to manufacture powders, such as liquid metal atomization, chemical and thermal decomposition reactions, electrolyte deposition, sol-gel process and

Two kinds of millingcould be distinguished - conventional low energy milling ( e.g. in a ball mill) and high energy milling. In the last case, the most used devices are planetary ball mill, SPEX ball mill and attritor ball mill. High energy milling (HEM) and mechanical alloying (MA) are terms utilized to name the synthesis of alloys and compounds by means of intensive milling that causes mixing at an intimate level and deformation of the materials microstructure (Costa et al. as cited by Benjamin, 1976 and Suryanarayana, 2001). MA is a solid-state powder processing technique involving repeated welding, fracturing, and rewelding of powder particles in a high-energy ball mill. Originally developed to produce oxide-dispersion strengthened (ODS) nickel- and iron-base superalloys for applications in the aerospace industry, MA has now been shown to be capable of synthesizing a variety of equilibrium and non-equilibrium alloy phases starting from blended elemental or prealloyed powders. The non-equilibrium phases synthesized include supersaturated solid solutions, metastable crystalline and quasicrystalline phases, nanostructures, and amorphous alloys. Recent advances in these areas and also on disordering of ordered

**3. Powder Metallurgy (P/M) used to obtain composite materials**

intermetallics and mechanochemical synthesis of materials have been critically reviewed after discussing the process and process variables involved in MA (Suryanarayana, 2001). The process of MA consists of loading the powder mix and the grinding medium (generally hardened steel or tungsten carbide balls) in a container sealed usually under a protective atmosphere (to avoid/minimize oxidation and nitridation during milling) and milling for the desired length of time. There are some parameters that must be controlled in the milling process, such as ball to powder ratio, length of milling time, atmosphere and ball sizes (Suryanarayana, 2001).

Nowadays, mechanical alloying has been successfully applied to prepare Ni-based, Febased, and Al-based oxide dispersion strengthened (ODS) alloys and carbide dispersion strengthened (CDS) alloys. During preparation of these materials, the strengthened phase is directly put into the matrix powders as a raw material. After a long time of ball milling, mechanical alloying can be achieved to some extent, but it is unlikely to produce the optimum combination between the strengthened phase, and the matrix. Element powders which can form oxide or carbide are mixed with the matrix powders and then subjected to ball milling, the supersaturated solid solution will form first, from which the dispersion strengthened alloy can be fabricated by a precipitation of second phase during heat treatment, or by mechanical-chemical reaction. This is a new method to produce CDS and ODS alloys by mechanical alloying, which can be used to optimize the strengthened action of the second phase to the maximum (Wang et al., 2001). During mechanical alloying, the mechanical energy transfers from the balls to the alloy powders, which results in the severe plastic deformation and the introduction of dislocations, lattice distortion and vacancies. When hard particles are incorporated into the matrix they promote an effect of hardening.

#### **5. Mechanism of reinforcement**

The characteristics of metal matrix composite materials are determined by their microstructure and internal interfaces, which are affected by their production and thermal mechanical prehistory. The microstructure covers the structure of the matrix and the reinforced phase. The chemical composition, grain and/or sub-grain size, texture, precipitation behavior and lattice defects are of importance to the matrix. The second phase is characterized by its volume percentage, its kind, size, distribution and orientation. Locally varying internal tension - due to the different thermal expansion behavior of the two phases - is an additional influencing factor. With knowledge of the characteristics of the components, the volume percentages, the distribution and orientation, it might be possible to estimate the characteristics of metallic composite materials. The approximations usually proceed from ideal conditions, i.e. optimal boundary surface formation, ideal distribution (very small number of contacts of the reinforcements among themselves) and no influence of the component on the matrix (comparable structures and precipitation behavior). However, in reality a strong interaction arises between the components involved, so that these models can only indicate the potential of a material. The different micro-, macro- and meso-scaled models proceed from different conditions and are differently developed (Kainer, 2006 as cited by Clyner & Withers, 1993 and McDaniels et al., 1965).

Effect of the Additives of Nanosized Nb and Ta Carbides

(MCS = 13,78nm) present similar medium crystallite size (MCS).

**7.2 Results**

**Powder**

on Microstructure and Properties of Sintered Stainless Steel 239

Figures 1, 2, 3, 4, 5 and 6 show the x-ray diffraction patterns and micrographs of NbC and TaC used in this work. Only carbides characteristics peaks were identified. The carbides produced by UFRN show wider peaks than supplied by others due to the crystal refinement in accordance with previous work (Souza et al., 1999 and Medeiros, 2002). The carbides synthesized by UFRN are finer than those supplied by other manufacture. The small crystallite size and lattice strain values, provided by Rietveld's Method presented in the table 1, confirm it. Both carbides synthesized in UFRN, NbC (MCS = 16.67nm) and TaC

**Size (nm)**

**NbC (2**θ**)**

34.90 40.49 58.40 69.90 73.40 87.48 97.70 MCS

**NbC-U** 17.65 17.64 17.10 16.43 16.22 15.27 16.38 16.67 0.0009

**NbC-C** 75.77 75.65 77.22 79.44 80.61 84.73 80.35 79.11 0.000145

**TaC (2**θ**)**

**34.88 40.48 58.60 70.06 73.66 87.61 98.03 MCS**

**TaC-U** 15.91 15.47 14.14 13.39 13.18 12.43 11.97 13.78 0.00156

**TaC-C** 42.68 41.62 40.19 39.79 39.73 40.04 40.61 40.66 0.000102

MCS = Crystallite medium size; NbC-U and TaC-U = carbides synthesized by UFRN; NbC-C and TaC-C = carbides supplied by Johnson Matthey Company and Sigma Aldrich.

Table 1. Crystallite size and lattice strain obtained for the carbides.

**Lattice Strain (%)**

**Lattice Strain (%)**

(nm)

**(nm)**

#### **6. Effect of ceramic particles in a metal matrix**

Ceramic particles are some of the most widely used materials for reinforcing metal matrix. In this way, the aging process can accelerate in aging alloys and the hardness can be improved considerably with direct impact and wear behavior (Torralba, 2003). Plastic deformation of a metallic material is associated to sliding of crystalline planes of crystalline structure and, therefore, there is a hardening and increase of mechanical resistance of the material when obstacles are placed to this movement. The particles of the reinforcement distributed uniformly into the grain serve as the obstacles to displacement of the discordances. The mechanisms of hardening occur when this discordance goes through the reinforcements or surrounds these reinforcements to continue its movement. In this sense, these mechanisms promote the hardening of the material during the plastic deformation. Obviously, the closer the particles are, the better tensile strength is improved. This is due the difficulty in movement of the discordances through or around the particles.

Compared with their unreinforced alloys, steel matrix composites offer higher hardness and wear resistance and their elastic modulus is the highest among all machinable materials and materials which can be hardened. On the other hand, as in all metal matrix composites, reductions in the fracture properties (ductility, toughness) are to be expected. Finally, the incorporation of dispersed ceramics particles inside the matrix can improve sintering, corrosion and oxidation properties of composites.

#### **7. Case study – Effect of the dispersion of nanosized carbides (NbC-TaC) in the sintered microstructure of the stainless steel 316L**

Composites containing different hard phases have already been studied (Tjong and Lau, 1999). Hard phases added to the steel matrix were mainly refractory carbides such as NbC, TaC and VC (Gordo et al., 2000). WC-Co (Ban and Shaw, 2002), Cr2Al, TiCr2, SiC and VC (Abenojar et al., 2002 and Abenojar et al., 2003) were also investigated as a potential reinforcement of the steel matrixes. This case study proposes a practical analysis to improve density and hardness of hard particles, used as additive. Recently a possibility to produce very homogeneous compacts with relative density among 95% without exaggerated grain growth was reported. These results were attributed to the incorporation of nanosized carbides' particles into stainless steel matrixes (Gomes et al., 2007).

#### **7.1 Experimental procedure**

This case study was carried out with starting powders of water-atomized stainless steel 316L manufactured by Höganäs. Pure sample and samples with addition of up to 3% of NbC and with 3% of TaC produced in the UFRN (Federal University of Rio Grande do Norte, Brazil) laboratory, previously characterized were mechanically milled in a conventional ball mill for 24 hours and axially cold pressed in a cylindrical steel die at 700 MPa. Sintering was carried out in vacuum. Samples were heated up to 1290°C with heating rate 20°C/min and isothermally held for 30 and 60 minutes. Sintered samples were characterized by x-ray diffraction, scanning electron microscopy and density and hardness were measured.

#### **7.2 Results**

238 Sintering – Methods and Products

Ceramic particles are some of the most widely used materials for reinforcing metal matrix. In this way, the aging process can accelerate in aging alloys and the hardness can be improved considerably with direct impact and wear behavior (Torralba, 2003). Plastic deformation of a metallic material is associated to sliding of crystalline planes of crystalline structure and, therefore, there is a hardening and increase of mechanical resistance of the material when obstacles are placed to this movement. The particles of the reinforcement distributed uniformly into the grain serve as the obstacles to displacement of the discordances. The mechanisms of hardening occur when this discordance goes through the reinforcements or surrounds these reinforcements to continue its movement. In this sense, these mechanisms promote the hardening of the material during the plastic deformation. Obviously, the closer the particles are, the better tensile strength is improved. This is due the difficulty in movement of the discordances through or around

Compared with their unreinforced alloys, steel matrix composites offer higher hardness and wear resistance and their elastic modulus is the highest among all machinable materials and materials which can be hardened. On the other hand, as in all metal matrix composites, reductions in the fracture properties (ductility, toughness) are to be expected. Finally, the incorporation of dispersed ceramics particles inside the matrix can improve sintering,

**7. Case study – Effect of the dispersion of nanosized carbides (NbC-TaC) in**

Composites containing different hard phases have already been studied (Tjong and Lau, 1999). Hard phases added to the steel matrix were mainly refractory carbides such as NbC, TaC and VC (Gordo et al., 2000). WC-Co (Ban and Shaw, 2002), Cr2Al, TiCr2, SiC and VC (Abenojar et al., 2002 and Abenojar et al., 2003) were also investigated as a potential reinforcement of the steel matrixes. This case study proposes a practical analysis to improve density and hardness of hard particles, used as additive. Recently a possibility to produce very homogeneous compacts with relative density among 95% without exaggerated grain growth was reported. These results were attributed to the incorporation of nanosized

This case study was carried out with starting powders of water-atomized stainless steel 316L manufactured by Höganäs. Pure sample and samples with addition of up to 3% of NbC and with 3% of TaC produced in the UFRN (Federal University of Rio Grande do Norte, Brazil) laboratory, previously characterized were mechanically milled in a conventional ball mill for 24 hours and axially cold pressed in a cylindrical steel die at 700 MPa. Sintering was carried out in vacuum. Samples were heated up to 1290°C with heating rate 20°C/min and isothermally held for 30 and 60 minutes. Sintered samples were characterized by x-ray diffraction, scanning electron microscopy and density and

**6. Effect of ceramic particles in a metal matrix**

corrosion and oxidation properties of composites.

**7.1 Experimental procedure**

hardness were measured.

**the sintered microstructure of the stainless steel 316L**

carbides' particles into stainless steel matrixes (Gomes et al., 2007).

the particles.

Figures 1, 2, 3, 4, 5 and 6 show the x-ray diffraction patterns and micrographs of NbC and TaC used in this work. Only carbides characteristics peaks were identified. The carbides produced by UFRN show wider peaks than supplied by others due to the crystal refinement in accordance with previous work (Souza et al., 1999 and Medeiros, 2002). The carbides synthesized by UFRN are finer than those supplied by other manufacture. The small crystallite size and lattice strain values, provided by Rietveld's Method presented in the table 1, confirm it. Both carbides synthesized in UFRN, NbC (MCS = 16.67nm) and TaC (MCS = 13,78nm) present similar medium crystallite size (MCS).

MCS = Crystallite medium size; NbC-U and TaC-U = carbides synthesized by UFRN; NbC-C and TaC-C = carbides supplied by Johnson Matthey Company and Sigma Aldrich.

Table 1. Crystallite size and lattice strain obtained for the carbides.

Effect of the Additives of Nanosized Nb and Ta Carbides

on Microstructure and Properties of Sintered Stainless Steel 241

Fig. 3. Scanning electron micrographs of NbC as supplied from UFRN – 500x.

Fig. 4. Scanning electron micrograph of NbC as supplied from UFRN – 5000x.

Fig. 1. X-ray diffraction patterns of Niobium Carbide produced at UFRN and from Johnson Matthey Company.

Fig. 2. X-ray diffraction patterns of Tantalum Carbide produced at UFRN and from Sigma Aldrich.

Fig. 1. X-ray diffraction patterns of Niobium Carbide produced at UFRN and from Johnson

Fig. 2. X-ray diffraction patterns of Tantalum Carbide produced at UFRN and from Sigma

Matthey Company.

Aldrich.

Fig. 3. Scanning electron micrographs of NbC as supplied from UFRN – 500x.

Fig. 4. Scanning electron micrograph of NbC as supplied from UFRN – 5000x.

Effect of the Additives of Nanosized Nb and Ta Carbides

from UFRN – 1000x.

from UFRN – 1000x.

on Microstructure and Properties of Sintered Stainless Steel 243

Figures 7 and 8 show the micrographs of composite powders milled mechanically 24 hours, with addition (3%) of NbC (Fig. 7) and TaC (Fig. 8). The arrows indicate the carbide particles.

Fig. 7. Scanning electron micrographs of stainless steel powders reinforced with 3% of NbC

Fig. 8. Scanning electron micrographs of stainless steel powders reinforced with 3% of TaC

Fig. 5. Scanning electron micrographs of TaC as supplied from UFRN – 500x.

Fig. 6. Scanning electron micrographs of TaC as supplied from UFRN – 5000x.

Fig. 5. Scanning electron micrographs of TaC as supplied from UFRN – 500x.

Fig. 6. Scanning electron micrographs of TaC as supplied from UFRN – 5000x.

Figures 7 and 8 show the micrographs of composite powders milled mechanically 24 hours, with addition (3%) of NbC (Fig. 7) and TaC (Fig. 8). The arrows indicate the carbide particles.

Fig. 7. Scanning electron micrographs of stainless steel powders reinforced with 3% of NbC from UFRN – 1000x.

Fig. 8. Scanning electron micrographs of stainless steel powders reinforced with 3% of TaC from UFRN – 1000x.

Effect of the Additives of Nanosized Nb and Ta Carbides

sintered for 30 minutes and 60 minutes, respectively.

on Microstructure and Properties of Sintered Stainless Steel 245

Figures 11 and 12 present the scanning electron micrographs of samples with TaC-U and

Fig. 11. Scanning electron micrographs of 316L stainless steel reinforced with 3% of TaC

Fig. 12. Scanning electron micrographs of 316L stainless steel reinforced with 3% of TaC

from UFRN sintered at 1290°C with isotherm of 30 minutes.

from UFRN sintered at 1290°C with isotherm of 60 minutes.

Figures 9 and 10 present the scanning electron micrographs of samples with NbC-U and sintered for 30 minutes and 60 minutes, respectively.

Fig. 9. Scanning electron micrographs of 316L stainless steel reinforced with 3% of NbC from UFRN sintered at 1290°C with isotherm of 30 minutes.

Fig. 10. Scanning electron micrographs of 316L stainless steel reinforced with 3% of NbC from UFRN sintered at 1290°C with isotherm of 60 minutes.

Figures 9 and 10 present the scanning electron micrographs of samples with NbC-U and

Fig. 9. Scanning electron micrographs of 316L stainless steel reinforced with 3% of NbC from

Fig. 10. Scanning electron micrographs of 316L stainless steel reinforced with 3% of NbC

from UFRN sintered at 1290°C with isotherm of 60 minutes.

sintered for 30 minutes and 60 minutes, respectively.

UFRN sintered at 1290°C with isotherm of 30 minutes.

Figures 11 and 12 present the scanning electron micrographs of samples with TaC-U and sintered for 30 minutes and 60 minutes, respectively.

Fig. 11. Scanning electron micrographs of 316L stainless steel reinforced with 3% of TaC from UFRN sintered at 1290°C with isotherm of 30 minutes.

Fig. 12. Scanning electron micrographs of 316L stainless steel reinforced with 3% of TaC from UFRN sintered at 1290°C with isotherm of 60 minutes.

Effect of the Additives of Nanosized Nb and Ta Carbides

Engineering A. Vol. 335. pp. 1-5.

University Press. Cambridge.

Hard Materials.

31360-5.

pp. 636-642

University of Rio Grande do Norte.

**9. References**

79.

on Microstructure and Properties of Sintered Stainless Steel 247

Abenojar, J. ; Velasco, F. ; Torralba, J.A. ; Calero, J.A. & Marcè, R. (2002). Reinforcing 316L

Abenojar, J. ; Velasco, F. ; Bautista, A. ; Campos, M. ; Bas, J.A. & Torralba, J.M. (2003).

Ban, Z.G. & Shaw, L.L. (2002). Synthesis and processing of nanostructured WC-Co

Breval, E. (1995). Synthesis route to metal matrix composites with specific properties : a

Clyne, T.W. & Withers, P.J. (1993). An introduction to metal matrix composites. Cambridge

Gomes, U.U. (1995). Tecnologia dos pós – Fundamentos e aplicações. Vol. 1. Federal

Gomes, U.U. ; Furukava, M. ; Soares, S.R.S. ; Oliveira, L.A. & Souza, C.P. (2007). Carbide

Gordo, E. ; Velasco, F. ; Atón, N. & Torralba, J.M. (2000). Wear mechanisms in high speed steel reinforced with (NbC)p and (TaC)p MMCs. Wear. Vol. 239. pp. 251-259. Costa, F.A. ; Silva, A.G.P. ; Silva Jr, J.F. & Gomes, U.U. (2008). Composite Ta-Cu powders

Kainer, K.U. (2006). Metal matrix composites – custom-made materials for automotive and

McDanels, D.L. ; Jech, R.W. & Weeton, J.W. (1965). Analysis of stress-strain behavior of

Medeiros, F.F.P. (2002). Synthesis of tungstein of tungsten and niobium carbides at low

Pagounis, E. & Kindros, V.K. (1998). Processing and properties of particles reinforced steel matrix composites. Materials Science and Engineering A. Vol. 246. pp. 221-234. Ralph, B. ; Yuen, H.C. & Lee, W.B. (1997). The processing of metal matrix composites – an overview. Journal of Materials Processing Technology. Vol. 64. pp. 339-353. Souza, C.P. ; Favotto, C. ; Satre, P. ; Honoré, A. & Roubin, M. (1999). Preparation of tantalum

Suryanarayana, C. (2001). Mechanical alloying and milling. Progress in Materials Science.

Suryanarayana, C. ; Ivanov, E. & Boldyred, V.V. (2001). The science and technology of mechanical alloyng. Materials Science & Engineering A. Vol. 304-306. pp. 151-158.

distribution effect on sintering and mechanical properties of stainless steel 316L.

prepared by high energy milling. International Journal of Refractory Metals and

aerospace engineering. Wiley-VCH Velargh GmbH & Co, Weiheim. ISBN : 3-527-

tungsten-fiber-reinforced copper composites. Trans. Metall. Soc. AIME. Vol. 223.

temperature, through gas-solid reaction in fixed bed reactor. Doctoral Theses.

carbide from an organometallic precursor. Brazilian Journal of Chemical

materials. Journal of Materials Science. Vol. 37. pp 3397-3403. Benjamin, J.S. (1976). Mechanical alloying. Materials Science. Vol. 13. pp. 279-300.

review. Composites Engineering . Vol. 5. pp. 1127-1133.

Materials – IV International Symposium. Portugal.

Federal University of Rio Grande do Norte, Brazil.

Engineering. V. 16. pp. 3397-3403.

Vol. 46. pp. 1-184.

stainless steel with intermetallic and carbide particles. Materials Science and

Atmosphere influence in sintering process of stainless steel matrix composites reinforced with hard particles. Composites Science and Technology. Vol. 63. pp. 69-

The table 2 presents values of densities and hardness. The carbides have influence on densification. Samples reinforced with NbC show higher densification. However, the great influence on the hardness values which increase from 76HV to 140 HV with addition of nanosized particles of NbC was observed.

Table 2. Hardness and densities values from sintered samples.

#### **7.3 Conclusions**

The effects of nanosized particles of refractory carbides (NbC and TaC) on the sintering mechanism, denseness and hardness of stainless steel were investigated. The action of hardening mechanisms owed differences in fine particles size (10-20nm) results in great increase in the hardness values of the samples reinforced with carbides, especially for NbC.

The reinforced samples present very homogeneous composite raising the relative density to 93% for TaC and 95% for NbC. Great influence on the hardness that increase from 76.0 HV up to 115HV and even to 140HV, for the samples with nanosized reinforcement, was observed, owing to dispersion and precipitation of fine carbides in metallic matrix. The addition of carbides increases significantly the hardness of the stainless steel due to grain size reduction of the metallic matrix during sintering. Thus, the hardness does not increase only as a function of the density but mostly due to nanosized particles segregation on grain boundaries in the sintered microstructure.

#### **8. Acknowledgment**

The authors wish to thank CAPES, CNPq, LMCME-UFRN and NEPGN-UFRN.

#### **9. References**

246 Sintering – Methods and Products

The table 2 presents values of densities and hardness. The carbides have influence on densification. Samples reinforced with NbC show higher densification. However, the great influence on the hardness values which increase from 76HV to 140 HV with addition of

316L pure 30 7.10±0.07 89.0 76.0±5.2

316L pure 60 7.23±0.02 91.0 71.0±5.0

316L+NbC-U 30 7.53±0.05 94.6 133.0±8.3

316L+NbC-U 60 7.57±0.05 95.0 140±8.2

316L+TaC-U 30 7.57±0.07 93.0 94.0±9.8

316L+TaC-U 60 7.58±0.06 93.0 115.0±9.9

The effects of nanosized particles of refractory carbides (NbC and TaC) on the sintering mechanism, denseness and hardness of stainless steel were investigated. The action of hardening mechanisms owed differences in fine particles size (10-20nm) results in great increase in the hardness values of the samples reinforced with carbides, especially for

The reinforced samples present very homogeneous composite raising the relative density to 93% for TaC and 95% for NbC. Great influence on the hardness that increase from 76.0 HV up to 115HV and even to 140HV, for the samples with nanosized reinforcement, was observed, owing to dispersion and precipitation of fine carbides in metallic matrix. The addition of carbides increases significantly the hardness of the stainless steel due to grain size reduction of the metallic matrix during sintering. Thus, the hardness does not increase only as a function of the density but mostly due to nanosized particles segregation on grain

The authors wish to thank CAPES, CNPq, LMCME-UFRN and NEPGN-UFRN.

**Density Hardness**

**(1000gf) (g/cm³) (%)**

nanosized particles of NbC was observed.

**Sample Sintering time**

**(min)**

Table 2. Hardness and densities values from sintered samples.

boundaries in the sintered microstructure.

**8. Acknowledgment**

**7.3 Conclusions**

NbC.

**12**

G. Bagliuk

*Ukraine*

**Properties and Structure of Sintered**

The iron-carbon alloys are the most wide-spread powder structural materials due to relative simplicity of technologies and accessibility of input materials for their fabrication (Dorofeev,

However, the need for increasing of sintered steels mechanical and operational characteristics level causes the need of their supplementary alloying. The basic alloying elements, which have found a wide application in a practice of powder metallurgy, are copper, nickel, manganese,

Herewith, the significant capacities for manufacturing of low-alloy steels, which operational characteristic often are not inferior of the level of the same for steels, produced with application of traditional alloying modes, makes it possible the use of boron as basic alloying element (Dudrova, 1993; Gülsoy, 2004; Napara-Volgina, 2011; Suzuki, 2002). In our previous articles it has been demonstrated that boron addition to the base material is very efficient way for improving basic mechanical properties of sintered steels as well positive

It is well known that boron activates sintering of iron and steels because of formation of persistent liquid phase. In addition, boron has high affinity to oxygen. Thereby while using boron as one of the alloying elements for production of sintered steels improved mechanical

Furthermore it is necessary to take into account that solubility of boron in the iron is exceedingly small, so its addition in powder steels will cause the number of difficulties, that make it necessary to optimize just but one boron composition in the sintered alloy, but the method of its addition in the initial powder mixture as well. The most often used method of boron addition is based on application of boron carbide powder as boron containing element. But sintering of Fe-B4C powder mixtures is attended with formation of derivative porosity in consequence of dissolution of boron carbide particles in a contact with iron matrix. To eliminate the noted effect as boron containing additive Fe-B-C system master

So, in order to evaluate the ability of boron to activate sintering of the base powder preforms with aim of increasing of the mechanical properties and wear resistance, study was carried

properties with excellent combination of hardness and ductility could be obtained.

chromium, vanadium and molybdenum (Zhang, 2004; Dorofeev, 1986).

effect on activation of sintering was observed too.

alloy can be used (Napara-Volgina, 2008; Xiu, 1999)..

**1. Introduction**

1986; Metals Handbook , 1978).

**Boron Containing Carbon Steels**

*Institute for Problems of Materials Science*

### **Properties and Structure of Sintered Boron Containing Carbon Steels**

G. Bagliuk *Institute for Problems of Materials Science Ukraine*

#### **1. Introduction**

248 Sintering – Methods and Products

Tjong, S.C. & Lau, K.C.J. (1999). Sliding wear of stainless steel matrix composite reinforced

Torralba, J.M. ; Costa, C.E. & Velasco, F. (2003). P/M Aluminum matrix composites : an overview. Journal of Materials Processing Technology. Vol. 133. pp. 203-206. Wang, C.G.; Qi, B.S.; Bai, Y.J.; Wu, J. & Yang, J.F. (2001). Dispersion strengthened alloy due

with TiB2 particles. Materials Letters. Vol. 41. pp. 153-158.

to the precipitation of carbide during mechanical alloying.

The iron-carbon alloys are the most wide-spread powder structural materials due to relative simplicity of technologies and accessibility of input materials for their fabrication (Dorofeev, 1986; Metals Handbook , 1978).

However, the need for increasing of sintered steels mechanical and operational characteristics level causes the need of their supplementary alloying. The basic alloying elements, which have found a wide application in a practice of powder metallurgy, are copper, nickel, manganese, chromium, vanadium and molybdenum (Zhang, 2004; Dorofeev, 1986).

Herewith, the significant capacities for manufacturing of low-alloy steels, which operational characteristic often are not inferior of the level of the same for steels, produced with application of traditional alloying modes, makes it possible the use of boron as basic alloying element (Dudrova, 1993; Gülsoy, 2004; Napara-Volgina, 2011; Suzuki, 2002). In our previous articles it has been demonstrated that boron addition to the base material is very efficient way for improving basic mechanical properties of sintered steels as well positive effect on activation of sintering was observed too.

It is well known that boron activates sintering of iron and steels because of formation of persistent liquid phase. In addition, boron has high affinity to oxygen. Thereby while using boron as one of the alloying elements for production of sintered steels improved mechanical properties with excellent combination of hardness and ductility could be obtained.

Furthermore it is necessary to take into account that solubility of boron in the iron is exceedingly small, so its addition in powder steels will cause the number of difficulties, that make it necessary to optimize just but one boron composition in the sintered alloy, but the method of its addition in the initial powder mixture as well. The most often used method of boron addition is based on application of boron carbide powder as boron containing element. But sintering of Fe-B4C powder mixtures is attended with formation of derivative porosity in consequence of dissolution of boron carbide particles in a contact with iron matrix. To eliminate the noted effect as boron containing additive Fe-B-C system master alloy can be used (Napara-Volgina, 2008; Xiu, 1999)..

So, in order to evaluate the ability of boron to activate sintering of the base powder preforms with aim of increasing of the mechanical properties and wear resistance, study was carried

Properties and Structure of Sintered Boron Containing Carbon Steels 251

(a) (b)

The unalloyed sintered Astaloy CrM has a homogeneous ferritic microstructure (fig. 3,a). The microhardness of the ferritic grains is 150-170 HV0,02. The representative microstructure of the boron alloyed sintered specimens is shown in figure 3,b. Boron causes hardening of the matrix, and its microhardness is very sensitive of the form of boron. For elemental boron with increasing boron content the microhardness of matrix ranges between 270-300 and 330-370 HV0,02 respectively. For B4C with increasing boron content the microhardness of matrix ranges 330-370 and 400-470 HV0.02 respectively. Beside, as a consequence of the eutectic reaction between austenite matrix and borides, the microstructure contains second constituent, which forms for higher boron content a continuous network on the grain boundaries. Figure 3,b shows the typical representative morphology of the second constituent, observed in 0.4 wt % boron (B4C) alloyed sintered specimen. Two different borides phases were found in the second constituent. The gray component is the Fe-Cr rich boride, while the white part of the component is the Fe-Mo

(a) (b)

Fig. 3. SEM micrograph of 0.4 (wt.) % boron (B4C) alloyed sintered Astaloy CrM.

Fig. 2. Optical micrograph of unetched boron-free (a) and with 0.4 (wt.) %

boron (B4C) alloyed sintered Astaloy CrM.

rich boride.

×500 (a); ×2500 (b).

on influence of boron composition and modes of its addition in the initial powder mixture on basic mechanical and tribotechnical characteristics of sintered steel.

#### **2. Effect of powder mixture content and modes of alloy addition on the structure and properties of sintered steels**

In (Kazior, 2002) the authors had studied the influence of boron addition in different forms (elemental boron powder, boron carbide B4C powder, and mixture of boron and carbide elemental powders) and different weight percentages on density and mechanical properties of prealloyed Astaloy CrM powders alloys.

The specimens were prepared by mixing of commercial water atomized Astaloy CrM powder with above-mentioned boron containing powders and were compacted at a pressure of 600 MPa to obtain an average density of 6.9 g/cm3. The green parts were sintered at 1250 0C for 30 min. in laboratory furnace in pure dry hydrogen.

The obtained results had shown, that the additive of boron caused densification, but the effect depends on the amount and form of boron additives. For 0,2 (wt.) % B only for boron carbide small increase in density can be observed, whereas for the other modes of boron addition density is of the same level as for boron-free sintered steel (fig. 1). A significant increase in sintered density was obtained with higher content of boron (0.4 wt. %), reaching 7.44 g/cm3 for iron–boron carbide powders mixture.

Fig. 1. Sintered density as a function of boron content and mode of its addition (Kazior, 2002).

Figure 2 shows the pore morphology of the boron-free and with 0.4 (wt.) % B4C sintered alloys. It is evident, that boron addition considerably changed the pore morphology: while the boron-free sintered alloy has in its structure the irregular pores, typical for solid phase sintering (fig. 2, a), otherwise the boron containing sample has the large well rounded pores, which are typically formed during liquid phase sintering (fig. 2, b).

on influence of boron composition and modes of its addition in the initial powder mixture

In (Kazior, 2002) the authors had studied the influence of boron addition in different forms (elemental boron powder, boron carbide B4C powder, and mixture of boron and carbide elemental powders) and different weight percentages on density and mechanical properties

The specimens were prepared by mixing of commercial water atomized Astaloy CrM powder with above-mentioned boron containing powders and were compacted at a pressure of 600 MPa to obtain an average density of 6.9 g/cm3. The green parts were

The obtained results had shown, that the additive of boron caused densification, but the effect depends on the amount and form of boron additives. For 0,2 (wt.) % B only for boron carbide small increase in density can be observed, whereas for the other modes of boron addition density is of the same level as for boron-free sintered steel (fig. 1). A significant increase in sintered density was obtained with higher content of boron (0.4 wt. %), reaching

Fig. 1. Sintered density as a function of boron content and mode of its addition (Kazior,

which are typically formed during liquid phase sintering (fig. 2, b).

Figure 2 shows the pore morphology of the boron-free and with 0.4 (wt.) % B4C sintered alloys. It is evident, that boron addition considerably changed the pore morphology: while the boron-free sintered alloy has in its structure the irregular pores, typical for solid phase sintering (fig. 2, a), otherwise the boron containing sample has the large well rounded pores,

**2. Effect of powder mixture content and modes of alloy addition on the**

on basic mechanical and tribotechnical characteristics of sintered steel.

sintered at 1250 0C for 30 min. in laboratory furnace in pure dry hydrogen.

**structure and properties of sintered steels**

7.44 g/cm3 for iron–boron carbide powders mixture.

2002).

of prealloyed Astaloy CrM powders alloys.

Fig. 2. Optical micrograph of unetched boron-free (a) and with 0.4 (wt.) % boron (B4C) alloyed sintered Astaloy CrM.

The unalloyed sintered Astaloy CrM has a homogeneous ferritic microstructure (fig. 3,a). The microhardness of the ferritic grains is 150-170 HV0,02. The representative microstructure of the boron alloyed sintered specimens is shown in figure 3,b. Boron causes hardening of the matrix, and its microhardness is very sensitive of the form of boron. For elemental boron with increasing boron content the microhardness of matrix ranges between 270-300 and 330-370 HV0,02 respectively. For B4C with increasing boron content the microhardness of matrix ranges 330-370 and 400-470 HV0.02 respectively. Beside, as a consequence of the eutectic reaction between austenite matrix and borides, the microstructure contains second constituent, which forms for higher boron content a continuous network on the grain boundaries. Figure 3,b shows the typical representative morphology of the second constituent, observed in 0.4 wt % boron (B4C) alloyed sintered specimen. Two different borides phases were found in the second constituent. The gray component is the Fe-Cr rich boride, while the white part of the component is the Fe-Mo rich boride.

Fig. 3. SEM micrograph of 0.4 (wt.) % boron (B4C) alloyed sintered Astaloy CrM. ×500 (a); ×2500 (b).

Properties and Structure of Sintered Boron Containing Carbon Steels 253

(a) (b)

(a) (b)

With increasing of sintering temperature to 1050 0C around B4C particle the new phase appears, which composition was identified by microX-ray spectrum analysis (fig. 6) as

The above presented results testify, that boron carbide is the compound, that is easily dissociates in a contact with iron at relatively low temperatures. Herewith B4C is a source of atomic boron and carbon, which are reacting with iron with establishing of new strengthening phases. Boron carbide is a sintering activator as well. However, thereupon that B4C is the compound relatively poor with carbon, the opportunity of supplementary

To evaluate both boron and carbon content in the initial powder mixture on structure and properties of sintered steel mixtures were prepared by mixing of water atomized pure iron powder with particles size smaller than 160 µm, 1÷3 % (weight) commercial boron carbide

Fig. 5. Optical micrograph of sintered at 900 0C for 10 (a) and 120 (b) min.

boron cementite with ∼75 % (atomic weight) Fe, ∼15 % C and ∼10 % B.

strengthening of Fe-B4C system by additive of carbon powder probably exists.

Fe+3 % (wt.) B4C powder mixture.

and 0,5÷2,0 % graphite powders.

Fig. 4. Fracture surface of unalloyed (a) and 0.4 % boron alloyed (b) sintered Astaloy CrM.

Table 1 summarizes the effect of boron on mechanical properties of the studied materials. The general trend is that, the higher boron content, the higher are the yield, strength, hardness and lower elongation of course. However, the results indicate that form of boron has significant effect on the tensile properties. For alloyed sintered specimens, both by boron carbide or elemental boron and graphite mixture, the microhardness of the matrix and in consequences tensile properties are higher as compared with elemental boron additive. Furthermore, it is clear that the tensile properties are not simply related to density. Boron causes both hardening of the matrix and densification, but the prevailing effect on properties is hardening. For example for 0.2 (wt.) % boron contents, despite of the increase in tensile strength, density does not increase significantly.

The above considerations are confirmed by fracture surface analyses. Figure 4 shows the representative fracture surface of tensile specimens of unalloyed and 0.4 (wt.) % boron (B4C) alloyed specimens, respectively. The fracture surface of the unalloyed material (fig. 3,a) presents the typical fracture morphology of a sintered material with interconnected porosity, while 0.4 (wt.) % boron (B4C) alloyed material fracture surface has a brittle character and crack propagation occurs along the eutectic constituent (fig. 3,b).

Table 1. Mechanical properties of boron-free and boron alloyed sintered Astaloy CrM (Kazior, 2002).

The research of structure formation at sintering of Fe - B4C powder mixtures had shown, that active interaction of the components takes place already at temperatures of 850÷900 0C with generation of Fe2B and FeB (at ∼1100 0C) phases (Turov, 1991). During heating of powder mixture yet before the beginning of diffusion of the boron to iron matrix contact surface of B4C particle with iron diminishes owing to considerable difference of their coefficients of thermal expansion (fig. 5,a). Despite of that, at 900÷950 0C it is beginning the formation of the new phases in neighborhood of boron carbide particles (fig. 5,b). This phase has a typical for boride layers needle-shaped structure. The X-ray analysis of insoluble residue, which was isolated by means of electrolytic etching, has shown the presence in a structure of the material of FeB and Fe2B phases with microhardness of 14÷18 GPa.

Table 1 summarizes the effect of boron on mechanical properties of the studied materials. The general trend is that, the higher boron content, the higher are the yield, strength, hardness and lower elongation of course. However, the results indicate that form of boron has significant effect on the tensile properties. For alloyed sintered specimens, both by boron carbide or elemental boron and graphite mixture, the microhardness of the matrix and in consequences tensile properties are higher as compared with elemental boron additive. Furthermore, it is clear that the tensile properties are not simply related to density. Boron causes both hardening of the matrix and densification, but the prevailing effect on properties is hardening. For example for 0.2 (wt.) % boron contents, despite of the increase

The above considerations are confirmed by fracture surface analyses. Figure 4 shows the representative fracture surface of tensile specimens of unalloyed and 0.4 (wt.) % boron (B4C) alloyed specimens, respectively. The fracture surface of the unalloyed material (fig. 3,a) presents the typical fracture morphology of a sintered material with interconnected porosity, while 0.4 (wt.) % boron (B4C) alloyed material fracture surface has a brittle

Astaloy CrM 157 287 4.36 182 Astaloy CrM + 0.2 % B 463 613 1.99 187 Astaloy CrM + 0.4 % B 515 688 1.43 247

(B4C) 509 666 1.74 272

(B4C) 592 741 1.27 302

(B4+C) 515 725 1.76 268

(B4+C) 662 806 1.57 320

The research of structure formation at sintering of Fe - B4C powder mixtures had shown, that active interaction of the components takes place already at temperatures of 850÷900 0C with generation of Fe2B and FeB (at ∼1100 0C) phases (Turov, 1991). During heating of powder mixture yet before the beginning of diffusion of the boron to iron matrix contact surface of B4C particle with iron diminishes owing to considerable difference of their coefficients of thermal expansion (fig. 5,a). Despite of that, at 900÷950 0C it is beginning the formation of the new phases in neighborhood of boron carbide particles (fig. 5,b). This phase has a typical for boride layers needle-shaped structure. The X-ray analysis of insoluble residue, which was isolated by means of electrolytic etching, has shown the presence in a

Table 1. Mechanical properties of boron-free and boron alloyed sintered Astaloy CrM

structure of the material of FeB and Fe2B phases with microhardness of 14÷18 GPa.

MTS (MPa) Elongation (%)

Hardness HV 10

character and crack propagation occurs along the eutectic constituent (fig. 3,b).

stress (MPa)

in tensile strength, density does not increase significantly.

Material 0.2 % offset yield

Astaloy CrM + 0.2 % B

Astaloy CrM + 0.4 % B

Astaloy CrM + 0.2 % B

Astaloy CrM + 0.4 % B

(Kazior, 2002).

Fig. 4. Fracture surface of unalloyed (a) and 0.4 % boron alloyed (b) sintered Astaloy CrM.

Fig. 5. Optical micrograph of sintered at 900 0C for 10 (a) and 120 (b) min. Fe+3 % (wt.) B4C powder mixture.

With increasing of sintering temperature to 1050 0C around B4C particle the new phase appears, which composition was identified by microX-ray spectrum analysis (fig. 6) as boron cementite with ∼75 % (atomic weight) Fe, ∼15 % C and ∼10 % B.

The above presented results testify, that boron carbide is the compound, that is easily dissociates in a contact with iron at relatively low temperatures. Herewith B4C is a source of atomic boron and carbon, which are reacting with iron with establishing of new strengthening phases. Boron carbide is a sintering activator as well. However, thereupon that B4C is the compound relatively poor with carbon, the opportunity of supplementary strengthening of Fe-B4C system by additive of carbon powder probably exists.

To evaluate both boron and carbon content in the initial powder mixture on structure and properties of sintered steel mixtures were prepared by mixing of water atomized pure iron powder with particles size smaller than 160 µm, 1÷3 % (weight) commercial boron carbide and 0,5÷2,0 % graphite powders.

Properties and Structure of Sintered Boron Containing Carbon Steels 255

Fe-3 % B4C, which hardness for mixtures with 1.5 and 2.0 % additive graphite considerably decreases (fig. 7,b, plot 3) owing to respective increase of specimens porosity (fig. 7,a).

The investigation of structure of sintered at relatively low temperatures (1100 0C) materials had shown, that for all steel compositions on a basis of ferrite-pearlite matrix structures contain multiple isolated inclusions, located in pores, that correspond to the configuration of boron carbide particles (fig. 8,a). Herewith, quantity of such inclusions increases with rise of

(a) (b)

Rise of sintering temperature up to 1150÷1200 0C is attended by dissolution of carbon in the matrix, formation of Fe2B + γ eutectic and severable healing of pores with inclusions and, as a consequence of this effect, alteration of structure (fig. 8, c-f). After sintering at 1150 0C material structure consist of substantially pearlite with trace of abnormal inclusions and isolated eutectic deposits (fig. 8, c), while after 1200 0C - pearlite inclusions on a basis of

With rise of carbon content the microhardness and amount of eutectic in sintered material increase. So, at enhancing of graphite content in a mixture from 0.5 to 2.0 % the eutectic

After sintering of all examined materials at 1150 0С some free carbon remain in their structure; its amount rises from 0.07 to 0.39 % in proportion to graphite content in the initial mixture.

Thermal treatment of the sintered materials results in essential advance of hardness, which values for the steels with 1 % B4C acquire 35÷40 HRC and with 3 % B4C - 50÷56 HRC, while for the boron-free specimens – from 25÷27 to 40÷43 HRC (subjected to carbon content) (fig. 9).

microhardness increased from 7.4 to 9.2 GPa, while its amount – from 50 to 80 %.

Fig. 7. Effect of graphite content on porosity (a) and hardness (b) of sintered at 1150 0C

boron carbide content in the mixture.

steels: 1 – boron-free; 2 – 1% B4C; 3 – 3 % B4C.

eutectic and separate pearlite areas with cementite net (fig. 8,e).

Powders were compacted at a pressure of 800 MPa to obtain cylindrical specimens with ∅10 and h≈10÷12 mm. The green parts were sintered at 1100÷1200 0C for 60 min. in the container with fusible glass bath gate (Fedorchenko, 1972). To estimate the influence of carbon on properties of the material the same specimens were prepared with boron carbide but without graphite additive. After sintering some specimens were quenched from 1000 0C in a water with subsequent tempering at 250 0C.

Density, hardness and compression strength of all specimens were examined. For the purpose of determining changes of structure and identifying phases, structural investigations by optical microscopy and phases microhardness estimations were carried out.

Figure 7 shows the influence of graphite content in the powder Fe-B4C mixtures on porosity and hardness of sintered at 1150 0C materials. As it can be seen from fig. 7,a porosity of compositions rises with increase of graphite content, and very considerably for Fe-3 % B4C in consequence of worse powder compressibility at compaction of B4C rich mixtures.

Fig. 6. Microstructure of sintered at 1050 0C Fe+3 % (wt.) B4C steel and elements distribution in the area of B4C particle.

Hardness of the sintered materials with 1 and 3 % of B4C with addition of graphite depends on two factors: values of their porosity and microstructure nature. So, with rise of B4C and graphite content hardness of sintered steels increase for all mixture compositions except of

Powders were compacted at a pressure of 800 MPa to obtain cylindrical specimens with ∅10 and h≈10÷12 mm. The green parts were sintered at 1100÷1200 0C for 60 min. in the container with fusible glass bath gate (Fedorchenko, 1972). To estimate the influence of carbon on properties of the material the same specimens were prepared with boron carbide but without graphite additive. After sintering some specimens were quenched from 1000 0C in a

Density, hardness and compression strength of all specimens were examined. For the purpose of determining changes of structure and identifying phases, structural investigations by optical

Figure 7 shows the influence of graphite content in the powder Fe-B4C mixtures on porosity and hardness of sintered at 1150 0C materials. As it can be seen from fig. 7,a porosity of compositions rises with increase of graphite content, and very considerably for Fe-3 % B4C

Fig. 6. Microstructure of sintered at 1050 0C Fe+3 % (wt.) B4C steel and elements distribution

Hardness of the sintered materials with 1 and 3 % of B4C with addition of graphite depends on two factors: values of their porosity and microstructure nature. So, with rise of B4C and graphite content hardness of sintered steels increase for all mixture compositions except of

in consequence of worse powder compressibility at compaction of B4C rich mixtures.

water with subsequent tempering at 250 0C.

in the area of B4C particle.

microscopy and phases microhardness estimations were carried out.

Fe-3 % B4C, which hardness for mixtures with 1.5 and 2.0 % additive graphite considerably decreases (fig. 7,b, plot 3) owing to respective increase of specimens porosity (fig. 7,a).

The investigation of structure of sintered at relatively low temperatures (1100 0C) materials had shown, that for all steel compositions on a basis of ferrite-pearlite matrix structures contain multiple isolated inclusions, located in pores, that correspond to the configuration of boron carbide particles (fig. 8,a). Herewith, quantity of such inclusions increases with rise of boron carbide content in the mixture.

Fig. 7. Effect of graphite content on porosity (a) and hardness (b) of sintered at 1150 0C steels: 1 – boron-free; 2 – 1% B4C; 3 – 3 % B4C.

Rise of sintering temperature up to 1150÷1200 0C is attended by dissolution of carbon in the matrix, formation of Fe2B + γ eutectic and severable healing of pores with inclusions and, as a consequence of this effect, alteration of structure (fig. 8, c-f). After sintering at 1150 0C material structure consist of substantially pearlite with trace of abnormal inclusions and isolated eutectic deposits (fig. 8, c), while after 1200 0C - pearlite inclusions on a basis of eutectic and separate pearlite areas with cementite net (fig. 8,e).

With rise of carbon content the microhardness and amount of eutectic in sintered material increase. So, at enhancing of graphite content in a mixture from 0.5 to 2.0 % the eutectic microhardness increased from 7.4 to 9.2 GPa, while its amount – from 50 to 80 %.

After sintering of all examined materials at 1150 0С some free carbon remain in their structure; its amount rises from 0.07 to 0.39 % in proportion to graphite content in the initial mixture.

Thermal treatment of the sintered materials results in essential advance of hardness, which values for the steels with 1 % B4C acquire 35÷40 HRC and with 3 % B4C - 50÷56 HRC, while for the boron-free specimens – from 25÷27 to 40÷43 HRC (subjected to carbon content) (fig. 9).

Properties and Structure of Sintered Boron Containing Carbon Steels 257

Fig. 9. Effect of graphite content on HRC hardness of sintered and thermal treated steels:

Fig. 10. Effect of graphite content on compression strength of sintered (1-4) and thermal treated (4a) steels with 1 % B4C (1), 3 % B4C (4, 4а) and boron-free Fe (2,3). Sintering

**3. Application of boron containing master alloys for manufacture of sintered**

As it can be seen from the above presented results, sintering of Fe-B4C powder mixtures is attended with formation of derivative porosity in consequence of dissolution of boron

1 – Fe + 3 % B4C; 2 – Fe + 1,5 % B4C; 3 – boron-free Fe.

temperature: 1, 2 – 1200 0С; 3, 4 – 1150 0С.

**steels**

Investigation of sintered and heat treated materials compression strength had shown, that ultimate values of durability have the steels with 1 % B4C and 1÷1.5 % additive carbon (fig. 10, plot 1). The least strength is for materials with 3 % B4C and carbon content more than 1.0 % (fig. 10, plots 4 and 4a). That kind of compositions are characterized by fragile nature of destruction, that can be explained by relatively high porosity (see fig. 7,a) and presence of significant amount of fragile eutectic.

Fig. 8. Structure of the materials sintered at 1100 (a, b); 1150 (c, d, f) and 1200 0С (e); composition of mixtures: Fe + 1 % B4C + 2 % C (a, c, e); Fe + 3% B4C + 2 % C (b, d, f).

Investigation of sintered and heat treated materials compression strength had shown, that ultimate values of durability have the steels with 1 % B4C and 1÷1.5 % additive carbon (fig. 10, plot 1). The least strength is for materials with 3 % B4C and carbon content more than 1.0 % (fig. 10, plots 4 and 4a). That kind of compositions are characterized by fragile nature of destruction, that can be explained by relatively high porosity (see fig. 7,a) and presence of

(a) (b)

(c) (d)

 (e) (f)

Fig. 8. Structure of the materials sintered at 1100 (a, b); 1150 (c, d, f) and 1200 0С (e); composition of mixtures: Fe + 1 % B4C + 2 % C (a, c, e); Fe + 3% B4C + 2 % C (b, d, f).

significant amount of fragile eutectic.

Fig. 9. Effect of graphite content on HRC hardness of sintered and thermal treated steels: 1 – Fe + 3 % B4C; 2 – Fe + 1,5 % B4C; 3 – boron-free Fe.

Fig. 10. Effect of graphite content on compression strength of sintered (1-4) and thermal treated (4a) steels with 1 % B4C (1), 3 % B4C (4, 4а) and boron-free Fe (2,3). Sintering temperature: 1, 2 – 1200 0С; 3, 4 – 1150 0С.

#### **3. Application of boron containing master alloys for manufacture of sintered steels**

As it can be seen from the above presented results, sintering of Fe-B4C powder mixtures is attended with formation of derivative porosity in consequence of dissolution of boron

Properties and Structure of Sintered Boron Containing Carbon Steels 259

Content of B4C in a mixture, %

0,4 2 6,8 - 0,17

0,8 2 13,6 - 0,35

1,2 2 20,3 - 0,52

1,6 2 27,1 - 0,69

1 11,0 - 0,19

3 4,6 - 0,16 4 - 0,5 0,10

1 22,2 - 0,37

3 9,2 - 0,32 4 - 1,0 0,20

1 33.3 - 0,56

3 13,7 - 0,48 4 - 1,5 0,30

1 44,4 - 0,76

3 18,4 - 0,64 4 - 2,0 0,40

Table 3. Composition of powder mixtures, used for manufacturing of boron containing

Fig. 11. Effect of boron content on porosity θ of green compacts (1, 2, 3) and sintered at 1200 <sup>0</sup>С preforms (1а, 2а, 3а) for different modes of his addition: В4С (1, 1а); amorphous boron (2,

Carbon content in a mixture, %

Boron content in a mixture, %

Mixture composition, %

Content of master alloy in a mixture, %

Master alloy

No. of master alloy (from table.1)

No.

I

II

III

IV

sintered steels.

2а); master alloy with 6% of B (3, 3а).

carbide particles in a contact with iron matrix. To eliminate the noted effect as boron containing additive Fe-B-C system master alloy can be used. The last one can increase the boron uniformity in the volume of sintered steels too, that probably result in improvement of their mechanical properties.

For estimation of the proposed approach several routes of boron addition to a powder mixture on the amount from 0,3 to 2,4 % and some mechanical properties of the sintered steels had been investigated. As boron-containing components boron carbide, amorphous boron of high purity and Fe-B-C system master alloy, produced by thermal synthesis from Fe and B4C powder mixtures, were used. All of the noted boron-containing components were mixed with water atomized iron powder with maximal particles size of 160 µm. Samples of ∅10×(10÷12) mm size were compacted by uniaxial die pressing at 800 MPa and sintered at 1100, 1150 and 1200 0С for 1,5 h in the container with fusible glass bath gate.

Three types of powder master alloys chemical compositions (table 2) were used for estimation of sintered steels mechanical properties. Boron content in the powder mixtures for the samples manufacturing was 0,4; 0,8; 1,2 and 1,6 % (weight) for every mode of boron addition (table 2).

Sintered iron samples with 0,5; 1,0; 1,5 and 2,0% boron carbide additives, that provided the same total boron in powder mixtures, were used as the reference ones.

Some sintered specimens were heat treated, i.e. water quenched from 1000 0С with subsequent tempering at 250 0C over a period of 3 hours.

Table 2. Chemical composition of master alloys and boron carbide, used for manufacturing of sintered steels.

Values of die pressed green compacts and sintered at 1200 0C preforms porosity for sintered steels, produced from different raw materials (powder mixtures with master alloy, B4C and amorphous boron powders) with different boron contents were investigated (fig. 11). It can be seen, that with increasing of boron concentration in the powder mixture preform green density decreased for all kinds of boron-containing components.

Maximal porosity values were observed for the green compacts, produced from master alloy containing powder mixtures (fig. 11, curve 3). It can be explained by significant its amount in the source mixture in comparison with other boron containing components and, owing to that, deterioration of such mixtures compressibility.

carbide particles in a contact with iron matrix. To eliminate the noted effect as boron containing additive Fe-B-C system master alloy can be used. The last one can increase the boron uniformity in the volume of sintered steels too, that probably result in improvement

For estimation of the proposed approach several routes of boron addition to a powder mixture on the amount from 0,3 to 2,4 % and some mechanical properties of the sintered steels had been investigated. As boron-containing components boron carbide, amorphous boron of high purity and Fe-B-C system master alloy, produced by thermal synthesis from Fe and B4C powder mixtures, were used. All of the noted boron-containing components were mixed with water atomized iron powder with maximal particles size of 160 µm. Samples of ∅10×(10÷12) mm size were compacted by uniaxial die pressing at 800 MPa and sintered at 1100, 1150 and 1200 0С for 1,5 h in the container with fusible glass bath

Three types of powder master alloys chemical compositions (table 2) were used for estimation of sintered steels mechanical properties. Boron content in the powder mixtures for the samples manufacturing was 0,4; 0,8; 1,2 and 1,6 % (weight) for every mode of boron

Sintered iron samples with 0,5; 1,0; 1,5 and 2,0% boron carbide additives, that provided the

Some sintered specimens were heat treated, i.e. water quenched from 1000 0С with

1 Master alloy 3.6 1.7 2 Master alloy 5.9 2.3 3 Master alloy 8.7 3.5 4 B4C 80.0 20.0

Table 2. Chemical composition of master alloys and boron carbide, used for manufacturing

Values of die pressed green compacts and sintered at 1200 0C preforms porosity for sintered steels, produced from different raw materials (powder mixtures with master alloy, B4C and amorphous boron powders) with different boron contents were investigated (fig. 11). It can be seen, that with increasing of boron concentration in the powder mixture preform green

Maximal porosity values were observed for the green compacts, produced from master alloy containing powder mixtures (fig. 11, curve 3). It can be explained by significant its amount in the source mixture in comparison with other boron containing components and, owing to

Boron content, % (weight)

Carbon content, % (weight)

same total boron in powder mixtures, were used as the reference ones.

subsequent tempering at 250 0C over a period of 3 hours.

density decreased for all kinds of boron-containing components.

that, deterioration of such mixtures compressibility.

boron addition

of their mechanical properties.

gate.

addition (table 2).

of sintered steels.

No. Modes of

Table 3. Composition of powder mixtures, used for manufacturing of boron containing sintered steels.

Fig. 11. Effect of boron content on porosity θ of green compacts (1, 2, 3) and sintered at 1200 <sup>0</sup>С preforms (1а, 2а, 3а) for different modes of his addition: В4С (1, 1а); amorphous boron (2, 2а); master alloy with 6% of B (3, 3а).

Properties and Structure of Sintered Boron Containing Carbon Steels 261

morphology is different: it is situated along the grain boundaries, shaping the developed framework (fig. 13, b), which can appear only as a result of liquid phase appearance in

According to X-ray micrography data, the composition of interface region, which appears around the boron carbide particle, is similar to boron cementite composition (about 75 % of

The feature of structure of sintered steels, manufactured from powder mixtures with boron carbide, unlike those of steels, made with use of master alloy, is the availability of diffusion

Mechanical properties testing results of sintered steels, produced from powder mixtures with different types of master alloys and boron carbide (table 2), had shown, that materials with 0.8-1.2 % B, produced from powder mixtures with low-alloy additives, provided higher values of hardness, than that, made with use of boron carbide as boron containing additives

porosity, generated owing to dissociation of boron carbide particles (fig. 14).

Fig. 13. Microstructure of steels with 1.0 % B4C sintered at 1050 (a) and 1200 0C (b).

(a) (b)

manufactured with application of master alloy (Fe - 5,3 % B - 2,3 % C) (a) and boron carbide

Fig. 14. Microstructure of sintered steels with 1,2 % B after sintering at 1150 0С,

after both sintering and thermal treatment (fig. 15, plots 1 and 2).

consequence of eutectic contact melting.

Fe, 15 % C and 10 % B).

(b); ×150.

However, after sintering at 1200 0С porosity of the samples, which contains master alloy admixtures, don't differ virtually from that, produced with use of the other modes of boron addition, since sintering of master alloy containing powders is accompanied by noticeable increased value of shrinkage. The denoted data are confirmed as well by phenomenon of shrinkage of the samples with master alloy for all sintering temperatures, unlike the materials with other boron containing components in the initial powder mixture, which are characterized by dilatation after sintering (fig. 12). The latter phenomenon is explained by the effect of boron carbide high temperature dissociation in contact with iron and high reaction activity of amorphous boron at process of sintering accompanied by formation of nonhealing "derivative" porosity in the areas of their location.

Fig. 12. Effect of boron content on linear shrinkage of billets after sintering at 1100 0С (a), 1150 0С (b) and 1200 0С (c) for different modes of boron insertion: В4С (1); amorphous boron (2); master alloy with 6 % of B (3).

After sintering of powder preforms with boron carbide, the last easily dissociates in contact with iron at relatively low temperatures and is a source of monatomic boron and carbon, which generate at interaction with Fe hard compounds, that reinforce the sintered metal. After sintering at 1050 0C it can be seen in the microstructure light-coloured and relatively hard (about 14 GPa) phase, located in the vicinity of incompletely dissociated boron carbide particles (fig. 13,a). Similar phase can be seen also after sintering at 1150-1200 0C, but its

However, after sintering at 1200 0С porosity of the samples, which contains master alloy admixtures, don't differ virtually from that, produced with use of the other modes of boron addition, since sintering of master alloy containing powders is accompanied by noticeable increased value of shrinkage. The denoted data are confirmed as well by phenomenon of shrinkage of the samples with master alloy for all sintering temperatures, unlike the materials with other boron containing components in the initial powder mixture, which are characterized by dilatation after sintering (fig. 12). The latter phenomenon is explained by the effect of boron carbide high temperature dissociation in contact with iron and high reaction activity of amorphous boron at process of sintering accompanied by formation of nonhealing "derivative" porosity in the areas of their

(a) (b)

(c) Fig. 12. Effect of boron content on linear shrinkage of billets after sintering at 1100 0С (a), 1150 0С (b) and 1200 0С (c) for different modes of boron insertion: В4С (1); amorphous boron

After sintering of powder preforms with boron carbide, the last easily dissociates in contact with iron at relatively low temperatures and is a source of monatomic boron and carbon, which generate at interaction with Fe hard compounds, that reinforce the sintered metal. After sintering at 1050 0C it can be seen in the microstructure light-coloured and relatively hard (about 14 GPa) phase, located in the vicinity of incompletely dissociated boron carbide particles (fig. 13,a). Similar phase can be seen also after sintering at 1150-1200 0C, but its

location.

(2); master alloy with 6 % of B (3).

morphology is different: it is situated along the grain boundaries, shaping the developed framework (fig. 13, b), which can appear only as a result of liquid phase appearance in consequence of eutectic contact melting.

According to X-ray micrography data, the composition of interface region, which appears around the boron carbide particle, is similar to boron cementite composition (about 75 % of Fe, 15 % C and 10 % B).

The feature of structure of sintered steels, manufactured from powder mixtures with boron carbide, unlike those of steels, made with use of master alloy, is the availability of diffusion porosity, generated owing to dissociation of boron carbide particles (fig. 14).

Mechanical properties testing results of sintered steels, produced from powder mixtures with different types of master alloys and boron carbide (table 2), had shown, that materials with 0.8-1.2 % B, produced from powder mixtures with low-alloy additives, provided higher values of hardness, than that, made with use of boron carbide as boron containing additives after both sintering and thermal treatment (fig. 15, plots 1 and 2).

Fig. 13. Microstructure of steels with 1.0 % B4C sintered at 1050 (a) and 1200 0C (b).

Fig. 14. Microstructure of sintered steels with 1,2 % B after sintering at 1150 0С, manufactured with application of master alloy (Fe - 5,3 % B - 2,3 % C) (a) and boron carbide (b); ×150.

Properties and Structure of Sintered Boron Containing Carbon Steels 263

So, for example, in a steel, sintered at 1150 0C with use of master alloy with 8.7 % of B, microhardness of eutectic is in the range of 6.3÷7.7 GPa, whereas the same for a steel,

Similar relationship can be seen too for microhardness of base material for steels, produced

Fig. 16. Effect of boron content on compression strength for steels sintered at 1080 0С,

The wear tests of the sintered steels with 1.2 % B (table 4) were performed in a pin on disc test bed in conditions of unlubricated friction with utilization of step loading method at loads 10-40 N/cm2 and sliding speed 5 m/s in air. As opposite element 65Mn steel with hardness 45 HRC had been used, while as a pattern material - 100Cr6 bearing steel (HRC 52). Evaluation for magnitude of wear was carried out by assessment of weight loss after

Wear testing results for three compositions of sintered steels are presented on fig. 17. It can be seen, that wear resistance for all kinds of sintered boron steels considerably exceeds that for 100Cr6 cast and rolled bearing steel. However, in spite of sufficiently near compositions of boron and carbon in the sintered steels under test, wear resistance of materials, manufactured with use of master alloys with increased composition of boron carbide (No. 2

manufactured with use of master alloys of different compositions

every stage of testing and translation the data into µm/km.

(No. 1–3, table 2) and boron carbide (4).

produced with master alloy with 5.3 % B − 8.58÷9.56 GPa.

with use of master alloys.

Similar behavior for relation of compression strength with boron content for different modes of boron addition in the powder mixtures (fig. 16) had been registered.

Indeed, hardness values of the same materials, produced with application of master alloy with 8,7 % B and boron carbide (fig. 15, plots 3 and 4) are considerably lower than that, produced with low-boron master alloy.

Fig. 15. Effect of boron content on hardness of sintered at 1150 0С (a) and heat-treated (b) steels, manufactured with use of master alloys of three compositions (No. 1–3, table 2) and boron carbide (4).

This difference in the mechanical properties of sintered steels with different kinds of boron source can be explained by more uniform distribution of low-boron containing elements in a balk of material, particularly – of eutectic constituent, in case of employment of master alloys in the source powder mixtures and higher content of carbon in the master alloy as compared with boron carbide.

The drawn conclusions were confirmed by the results of sintered steels metallographic study, which had shown, that microhardness of structural components in the materials, manufactured with use of master alloy with 8.7 % B4C (no. 3, table 2) is somewhat lower than that, produced with low-boron containing master alloys (No. 1 and 2) with 3.6 and 5.8 % of B.

Similar behavior for relation of compression strength with boron content for different modes

Indeed, hardness values of the same materials, produced with application of master alloy with 8,7 % B and boron carbide (fig. 15, plots 3 and 4) are considerably lower than that,

(a) (b)

This difference in the mechanical properties of sintered steels with different kinds of boron source can be explained by more uniform distribution of low-boron containing elements in a balk of material, particularly – of eutectic constituent, in case of employment of master alloys in the source powder mixtures and higher content of carbon in the master alloy as

The drawn conclusions were confirmed by the results of sintered steels metallographic study, which had shown, that microhardness of structural components in the materials, manufactured with use of master alloy with 8.7 % B4C (no. 3, table 2) is somewhat lower than that, produced with low-boron containing master alloys (No. 1 and 2) with 3.6 and 5.8

Fig. 15. Effect of boron content on hardness of sintered at 1150 0С (a) and heat-treated (b) steels, manufactured with use of master alloys of three compositions (No. 1–3, table 2)

of boron addition in the powder mixtures (fig. 16) had been registered.

produced with low-boron master alloy.

and boron carbide (4).

% of B.

compared with boron carbide.

So, for example, in a steel, sintered at 1150 0C with use of master alloy with 8.7 % of B, microhardness of eutectic is in the range of 6.3÷7.7 GPa, whereas the same for a steel, produced with master alloy with 5.3 % B − 8.58÷9.56 GPa.

Similar relationship can be seen too for microhardness of base material for steels, produced with use of master alloys.

Fig. 16. Effect of boron content on compression strength for steels sintered at 1080 0С, manufactured with use of master alloys of different compositions (No. 1–3, table 2) and boron carbide (4).

The wear tests of the sintered steels with 1.2 % B (table 4) were performed in a pin on disc test bed in conditions of unlubricated friction with utilization of step loading method at loads 10-40 N/cm2 and sliding speed 5 m/s in air. As opposite element 65Mn steel with hardness 45 HRC had been used, while as a pattern material - 100Cr6 bearing steel (HRC 52). Evaluation for magnitude of wear was carried out by assessment of weight loss after every stage of testing and translation the data into µm/km.

Wear testing results for three compositions of sintered steels are presented on fig. 17. It can be seen, that wear resistance for all kinds of sintered boron steels considerably exceeds that for 100Cr6 cast and rolled bearing steel. However, in spite of sufficiently near compositions of boron and carbon in the sintered steels under test, wear resistance of materials, manufactured with use of master alloys with increased composition of boron carbide (No. 2

Properties and Structure of Sintered Boron Containing Carbon Steels 265

The presented results had shown that application of boron containing master alloys provides a superior level of sintered and heat-treated boron steels mechanical properties as

Boron content in the sintered steels that provides an optimum combination of their properties with application of master alloys of different composition, is 1.2 % (weight). This content of B corresponds to ultimate values of hardness and strength for both sintered and

Comparison of the mechanical and tribotechnical characteristics of the tested materials testify that there is no direct relation among values of their strength, hardness and wear resistance properties. However, every of this steels, in spite of presence of residual porosity, while possessing by commensurable mechanical characteristics with 100Cr6 bearing rolled

Causton, R. J. (2001). Hardenability of sintered boron-carbon steels. *Metal Powder Report*, Vol.

Dorofeev, Yu. G. Marinenko, L. G. (1986). *Structural powder materials and production.*

Dudrova, E. et al. (1993). Properties of liquid phase sintered low alloy steel containing boron. *Pokroky Praskove Metalurgie.* No 3, 1993, pp. 63–74. (In Czech). Fedorchenko, I. M. et al. (1972). Technology of sintering metal-ceramic materials without use

Gülsoy, H.Ö. et al. (2004). Effect of FeB additions on sintering characteristics of injection

Kazior, J. et al. (2002). The influence of boron on the mechanical properties of prealloyed

Metals Handbook. Amer Society for metals. (1978). *Ferrous Powder Metallurgy Materials in*

Napara-Volgina, S. G. et al. (2008). Investigation of features of various modes of alloying by

Napara-Volgina, S. G. et al. (2011). Structure and Properties of Sintered Iron-Boron-Carbon

Suzuki, H. Y. et al (2002). Effect of boron on alloying and pore morphology in PM carbon steel. *J. Jpn. Soc. Powder Metal.*, 2002, vol. 49, no. 7, pp. 600-606. (In Japanese). Turov, Yu. V. et al. (1991). Structure formation at sintering of powder composition ironboron carbide. *Powder metallurgy.* 1991, no. 6, pp.25-31. (In Russian).

*Properties and Selection: Irons and steels*. Vol. I, 9th ed., 1978.

of protective flowing atmosphere. Powder metallurgy. 1972, no. 5, pp.26-32. (In

moulded 17-4PH stainless steel powder. *Journal of Materials Science*. 2004, August,

CrM powders. *Deformation and Fracture in Structural PM Materials. DF PM 2002*. Vol.

boron of powder structural steels. *Metal science and metal treatment.* 2008, no. 4, pp.

Alloys with Different Carbon Contents. *Powder Metallurgy and Metal Ceramics*, 2011,

heat-treated materials and high level of wear resistance as well.

steel, is distinguished by considerably raised wear resistance.

Metallurgia, Moscow, 1986. (In Russian).

56, Issue 6, June 2001, p.40.

vol. 39, no. 15, pp. 4835-4840.

**4. Conclusions**

**5. References**

Russian).

1, pp.125-131.

34-37. (In Ukrainian).

Vol. 50, No. 1-2, pp.67-72.

compared with boron carbide.

and 3) at moderate loads (50 and 75 N/cm2) considerably exceeds the same for the steels, synthesized with relatively "poor" master alloy (No. 1).

Table 4. Compositions of sintered and heat-treated boron steels for wear testing.

With increasing of loads to 100 N/cm2 wear characteristics for sintered steels of all compositions are flattening.

High wear resistance of boron sintered steels as compared with 100Cr6 bearing steel can be explained by their heterophase structure and, particularly, by presence of hard eutectic component, which amount growth with increase of boron and carbon composition in the source powder mixture.

Fig. 17. Effect of load on wear at unlubricated friction of sintered boron steels (No. 1-3, table 3) and 100Cr6 bearing rolled steel (4).

#### **4. Conclusions**

264 Sintering – Methods and Products

and 3) at moderate loads (50 and 75 N/cm2) considerably exceeds the same for the steels,

No. Composition of master alloy Composition of B and C in steels, %

1 3,6 % В + 1,7 % С 1,2 0,86 2,06 1,4 2 5,9 % В + 2,3 % С 1,2 0,82 2,02 1,5 3 8,7 % В + 3,5 % С 1,2 0,78 1,98 1,5

With increasing of loads to 100 N/cm2 wear characteristics for sintered steels of all

High wear resistance of boron sintered steels as compared with 100Cr6 bearing steel can be explained by their heterophase structure and, particularly, by presence of hard eutectic component, which amount growth with increase of boron and carbon composition in the

Fig. 17. Effect of load on wear at unlubricated friction of sintered boron steels (No. 1-3, table

Table 4. Compositions of sintered and heat-treated boron steels for wear testing.

В С В+С В/С

synthesized with relatively "poor" master alloy (No. 1).

compositions are flattening.

source powder mixture.

3) and 100Cr6 bearing rolled steel (4).

The presented results had shown that application of boron containing master alloys provides a superior level of sintered and heat-treated boron steels mechanical properties as compared with boron carbide.

Boron content in the sintered steels that provides an optimum combination of their properties with application of master alloys of different composition, is 1.2 % (weight). This content of B corresponds to ultimate values of hardness and strength for both sintered and heat-treated materials and high level of wear resistance as well.

Comparison of the mechanical and tribotechnical characteristics of the tested materials testify that there is no direct relation among values of their strength, hardness and wear resistance properties. However, every of this steels, in spite of presence of residual porosity, while possessing by commensurable mechanical characteristics with 100Cr6 bearing rolled steel, is distinguished by considerably raised wear resistance.

#### **5. References**

**13**

*1India 2Japan*

**Influence of Sintering Temperature**

**Nanocrystalline Manganites**

*Nagatsuta, Midori-ku, Yokohama,*

*1Department of Physics, Osmania University, Hyderabad,*

**on Magnetotransport Behavior of Some**

G. Venkataiah2, Y. Kalyana Lakshmi1 and P. Venugopal Reddy1,\*

The colossal magnetoresistance (CMR) in hole doped manganese oxides, widely known as manganites with formula Ln1-xAexMnO3 (where Ln is a trivalent rare earth and Ae is a divalent alkaline earth ion; x=0-1), has been intensively studied over the last two decades. In general, these systems comprise a strong competition between charge, orbital, lattice and spin degrees of freedom all of which make them an intriguing subject of research (Dogotto, 2003; Goodenough, 2003; Ramirez, 1997; Tokura, 2000). The end-members of this system are antiferromagnetic (AFI) insulators. Partial substitution of the Ln with divalent alkaline-earth ion of the ABO3 structure introduces mixed valence Mn4+/Mn3+ on the B site. With decreasing temperature the system undergoes a transition from the paramagnetic to the ferromagnetic state accompanied by a metal-insulator transition. A large number of studies on CMR materials of different forms such as single crystals (Okuda et al., 1998; Zhou et al., 1997), thin film (Kwon et al., 1997; Rao et al., 1998; Suzuki et al., 1997) and ceramics (Hwang et al., 1995; Mahendiran et al., 1996a, 1996b) for the basic research point of view and also the possible future device applications view point. The Zener (1951) has proposed a simplest model, known as double-exchange (DE) model to explain the electrical behavior in ferromagnetic metallic region below Curie Temperature (TC). However, due to various interactions among charge, spin and lattice makes these materials more complex and DE alone cannot explain the entire electrical transport behavior. Later on various theoretical models have been proposed by considering, electron-lattice and spin lattice interaction and even today there is no comprehensive model to explain transport phenomena in manganites

These doped perovskite manganites show large magnetoresistance (MR) around metal to insulator transition temperature (TP) at high magnetic fields and below room temperatures which restricts their applicability to hands on devices. Improved MR could be achieved for

**1. Introduction**

(Millis et al., 1995; Tokura, 2000).

Corresponding Author

 \*

*2Materials and Structures Laboratory, Tokyo Institute of Technology,*

### **Influence of Sintering Temperature on Magnetotransport Behavior of Some Nanocrystalline Manganites**

G. Venkataiah2, Y. Kalyana Lakshmi1 and P. Venugopal Reddy1,\* *1Department of Physics, Osmania University, Hyderabad, 2Materials and Structures Laboratory, Tokyo Institute of Technology, Nagatsuta, Midori-ku, Yokohama, 1India 2Japan*

#### **1. Introduction**

266 Sintering – Methods and Products

Zhang, Z. et al (2004). Mechanical properties of Fe–Mo–Mn–Si–C sintered steels. *Powder*

Xiu, Z.M., et al. (1999). Fe-Mo-B-C – sintered steels produced by addition of master alloy

powders. *Acta Met. Sin.* 1999. vol. 12, no. 5. – pp. 1198-1201.

*Metallurgy*, 2004, Vol. 47, No. 3, pp. 239-246.

The colossal magnetoresistance (CMR) in hole doped manganese oxides, widely known as manganites with formula Ln1-xAexMnO3 (where Ln is a trivalent rare earth and Ae is a divalent alkaline earth ion; x=0-1), has been intensively studied over the last two decades. In general, these systems comprise a strong competition between charge, orbital, lattice and spin degrees of freedom all of which make them an intriguing subject of research (Dogotto, 2003; Goodenough, 2003; Ramirez, 1997; Tokura, 2000). The end-members of this system are antiferromagnetic (AFI) insulators. Partial substitution of the Ln with divalent alkaline-earth ion of the ABO3 structure introduces mixed valence Mn4+/Mn3+ on the B site. With decreasing temperature the system undergoes a transition from the paramagnetic to the ferromagnetic state accompanied by a metal-insulator transition. A large number of studies on CMR materials of different forms such as single crystals (Okuda et al., 1998; Zhou et al., 1997), thin film (Kwon et al., 1997; Rao et al., 1998; Suzuki et al., 1997) and ceramics (Hwang et al., 1995; Mahendiran et al., 1996a, 1996b) for the basic research point of view and also the possible future device applications view point. The Zener (1951) has proposed a simplest model, known as double-exchange (DE) model to explain the electrical behavior in ferromagnetic metallic region below Curie Temperature (TC). However, due to various interactions among charge, spin and lattice makes these materials more complex and DE alone cannot explain the entire electrical transport behavior. Later on various theoretical models have been proposed by considering, electron-lattice and spin lattice interaction and even today there is no comprehensive model to explain transport phenomena in manganites (Millis et al., 1995; Tokura, 2000).

These doped perovskite manganites show large magnetoresistance (MR) around metal to insulator transition temperature (TP) at high magnetic fields and below room temperatures which restricts their applicability to hands on devices. Improved MR could be achieved for

<sup>\*</sup> Corresponding Author

Influence of Sintering Temperature

temperature range 77–300 K.

**0.67(Ln (NO3)3 ) solution**

**Ln = La, Nd**

**3. Results & discussions 3.1 Structural properties**

on Magnetotransport Behavior of Some Nanocrystalline Manganites 269

crucial factor in controlling the transport and magnetic properties. Mn4+ concentration has been determined by redox titrations (Vogel, 1978). The phase purity of the samples was checked by X-ray diffraction (XRD) and grain distribution was estimated by scanning electron microscopy (SEM). The electrical resistivity and magnetoresistance studies were undertaken by a Janis "supervaritemp" cryostat in applied elds of 0–7 T, over a temperature range of 77–300 K using a four-point probe method. a.c susceptibility measurements were carried out using a dual channel lock-in amplier, while thermoelectric power studies were carried out by a two-probe differential method over a

Fig. 1. A flow chart of Sol-gel method for the preparation of CMR manganites.

Most of the manganites crystallize in the various derivatives of the so-called perovskites structure named after the mineral perovskite, CaTiO3.Figure 2 illustrates the ideal cubic perovskite structure, with the general formula ABO3. The structure may be conceived as a close-packed array formed of O2− anions and A cations with B cations located at the

**Adjust PH value between 6.5 & 7 Evaporate to get thick viscous mass. Add**

**Add citric acid to the solution in 1:1 ratio**

**Thick viscous solution**

**Ash colored mass**

**Ln0.67Ae0.33MnO3**

**0.33 (Ae (NO3)2 )solution Ae = Ca, Sr, Ba, Na**

**Gel formation Evaporate slowly on hot plate**

**ethylene glycol in 1:1.2 ratios**

**Mn (NO3)3 solution**

**burn small amounts**

**Calcination followed by sintering at four different temperatures**

nanosized perovskite manganite samples prepared through the sol–gel process. The sol–gel process also has potential advantages over other traditional processing techniques such as better homogeneities, low processing temperature and improved material properties. Although there are several reports on the synthesis of nanoscale manganites, none of them were carried out systematically in the context of the influence of sintering temperature on the electrical, magnetic and magneto transport behavior of manganites. It is well know that when the size of the particles is reduced to a few nanometers some of the basic properties such as magnetoresistance, superparamagnetism, coercivity, Curie temperature and Saturation magnetization are effected when compared with their bulk counterparts. Several works on different nanoscale systems indicate that the surface is responsible for their apparent anomalous behavioral changes with reduced dimensions. In fact, the preparation procedure and sintering conditions determines the nature of the surface region of nanosize grains, which plays a very crucial role in electrical transport, magnetic, and magnetotransport behavior of nano dimensional systems. Therefore, it is felt that a detailed review on the influence of sintering on magnetotransport behavior of manganites is essential and important.

The senior author of the paper (P. V. Reddy) along with his group of students investigated the structural, magnetic, electrical and magnetotransport properties of a number doped manganites over a period of more than a decade and published a series of papers in a number of International Journals (Kalyana Lakshmi and Reddy 2008, 2009a, 2009b, 2010, 2011; Kalyana Lakshmi et al., 2009; Venkataiah et al., 2005, 2007a, 2007b, 2007c, 2008, 2010; Venkataiah and Reddy, 2005, 2009a, 2009b). A summary of the work based on the influence of sintering on various properties of the manganites substituted with both divalent and monovalent ions is presented here.

#### **2. Experimental procedure**

The samples with different particle sizes were prepared using sol-gel route followed by heat treatment at four different sintering temperatures (TS) (Venkataiah et al., 2005, 2007, 2008, 2010; Venkataiah and Reddy 2005, 2009a, 2009b; Kalyana Lakshmi and Reddy 2009a). Five series of samples i.e., La0.67Ca0.33MnO3, La0.67Sr0.33MnO3, La0.67Ba0.33MnO3, Nd0.67Sr0.33MnO3 and La0.67Na0.33MnO3 were prepared to study the effect of sintering temperature on various properties. The samples are hereafter designated as LCMO, LSMO, LBMO, NSMO and LNMO respectively and the samples sintered at 800, 900, 1000 and 1100 C are designated with 8, 9, 10 and 11 numerical representations followed by compositional abbreviations. The synthesis procedure of sol-gel method is outlined in Fig. 1.

The stoichiometric amounts of oxides / carbonates were taken and all of them were converted into metal nitrates and the resultant nitrate mixture solution was converted into citrate by adding 1:1 ratio of the citric acid to the metal atoms present in the composition. To carry out the reaction at greater speed, a proper reaction environment is to be created and hence pH of the solution was adjusted between 6.5 and 7 by adding ammonia solution. Finally, the powder was pressed into pellets and sintered in air at four different temperatures i.e., 800, 900, 1000 and 11000C for 4 hr. For a direct comparison of the properties of the samples as a function of sintering temperature (grain size) it is important to ensure that the Mn4+ concentration remains similar in different samples, since it is

nanosized perovskite manganite samples prepared through the sol–gel process. The sol–gel process also has potential advantages over other traditional processing techniques such as better homogeneities, low processing temperature and improved material properties. Although there are several reports on the synthesis of nanoscale manganites, none of them were carried out systematically in the context of the influence of sintering temperature on the electrical, magnetic and magneto transport behavior of manganites. It is well know that when the size of the particles is reduced to a few nanometers some of the basic properties such as magnetoresistance, superparamagnetism, coercivity, Curie temperature and Saturation magnetization are effected when compared with their bulk counterparts. Several works on different nanoscale systems indicate that the surface is responsible for their apparent anomalous behavioral changes with reduced dimensions. In fact, the preparation procedure and sintering conditions determines the nature of the surface region of nanosize grains, which plays a very crucial role in electrical transport, magnetic, and magnetotransport behavior of nano dimensional systems. Therefore, it is felt that a detailed review on the influence of sintering on magnetotransport behavior of manganites is

The senior author of the paper (P. V. Reddy) along with his group of students investigated the structural, magnetic, electrical and magnetotransport properties of a number doped manganites over a period of more than a decade and published a series of papers in a number of International Journals (Kalyana Lakshmi and Reddy 2008, 2009a, 2009b, 2010, 2011; Kalyana Lakshmi et al., 2009; Venkataiah et al., 2005, 2007a, 2007b, 2007c, 2008, 2010; Venkataiah and Reddy, 2005, 2009a, 2009b). A summary of the work based on the influence of sintering on various properties of the manganites substituted with both divalent and

The samples with different particle sizes were prepared using sol-gel route followed by heat treatment at four different sintering temperatures (TS) (Venkataiah et al., 2005, 2007, 2008, 2010; Venkataiah and Reddy 2005, 2009a, 2009b; Kalyana Lakshmi and Reddy 2009a). Five series of samples i.e., La0.67Ca0.33MnO3, La0.67Sr0.33MnO3, La0.67Ba0.33MnO3, Nd0.67Sr0.33MnO3 and La0.67Na0.33MnO3 were prepared to study the effect of sintering temperature on various properties. The samples are hereafter designated as LCMO, LSMO, LBMO, NSMO and LNMO respectively and the samples sintered at 800, 900, 1000 and 1100 C are designated with 8, 9, 10 and 11 numerical representations followed by compositional abbreviations. The

The stoichiometric amounts of oxides / carbonates were taken and all of them were converted into metal nitrates and the resultant nitrate mixture solution was converted into citrate by adding 1:1 ratio of the citric acid to the metal atoms present in the composition. To carry out the reaction at greater speed, a proper reaction environment is to be created and hence pH of the solution was adjusted between 6.5 and 7 by adding ammonia solution. Finally, the powder was pressed into pellets and sintered in air at four different temperatures i.e., 800, 900, 1000 and 11000C for 4 hr. For a direct comparison of the properties of the samples as a function of sintering temperature (grain size) it is important to ensure that the Mn4+ concentration remains similar in different samples, since it is

essential and important.

monovalent ions is presented here.

synthesis procedure of sol-gel method is outlined in Fig. 1.

**2. Experimental procedure**

crucial factor in controlling the transport and magnetic properties. Mn4+ concentration has been determined by redox titrations (Vogel, 1978). The phase purity of the samples was checked by X-ray diffraction (XRD) and grain distribution was estimated by scanning electron microscopy (SEM). The electrical resistivity and magnetoresistance studies were undertaken by a Janis "supervaritemp" cryostat in applied elds of 0–7 T, over a temperature range of 77–300 K using a four-point probe method. a.c susceptibility measurements were carried out using a dual channel lock-in amplier, while thermoelectric power studies were carried out by a two-probe differential method over a temperature range 77–300 K.

Fig. 1. A flow chart of Sol-gel method for the preparation of CMR manganites.

#### **3. Results & discussions**

#### **3.1 Structural properties**

Most of the manganites crystallize in the various derivatives of the so-called perovskites structure named after the mineral perovskite, CaTiO3.Figure 2 illustrates the ideal cubic perovskite structure, with the general formula ABO3. The structure may be conceived as a close-packed array formed of O2− anions and A cations with B cations located at the

Influence of Sintering Temperature

Intensity (a.u)

(a)

(b)

20 30 40 50 60 70 80

LSMO-8

LSMO-9 LSMO-10 LSMO-11

NSMO-8

NSMO-9 NSMO-10

NSMO-11

2 (Degree)

20 30 40 50 60 70 80

2 (Degree)

increases and porosity decreases.

on Magnetotransport Behavior of Some Nanocrystalline Manganites 271

Fig. 3. (a, b) XRD patterns of LSMO & NSMO group of manganites at various sintering

In this method, the experimental x-ray peak broadening is given by =tan +K/ st cos, where is the full width at half maximum (FWHM) of the XRD peaks after subtracting FWHM of the standard sample (SiO2), is the Bragg angle, K is the grain shape factor (0.89), is the wavelength of Cu K radiation, and is the strain in the sample. A plot between cos versus sin is shown in Figure 3(c). The average crystallite sizes were calculated from the intercept of the straight line with the vertical axis, while the slope of the line gives the lattice strain (). The average crystallite sizes are found to increase with increasing sintering temperature and are in nanometer range. The variation of crystallite size with sintering temperature is shown in Figure 4. In fact, the computed crystallite sizes are comparable with those obtained from SEM. Figure 5 shows the representative SEM images of LSMO samples sintered at 800 and 1100ºC. The SEM observation reveals that there is a uniform distribution of grain sizes for the samples and as the sintering temperature increases, the grain size

0.001

0.002

0.003

cos

0.004

0.005

**(c)**

0.006

0.2 0.3 0.4 0.5 0.6 0.7

sin

 **NSMO - 8 NSMO - 9 NSMO - 10 NSMO - 11**

temperatures (TS) (c) Williamson-Hall plot of Nd0.67Sr0.33MnO3 manganites.

octahedral interstitial sites. The BO6 octahedra make contact to each other by their vertices and form a three-dimensional network. In manganites the A-sites of the perovskite structure are occupied by trivalent rare earth ion while B-sites with Mn ions. The stability of the perovskite structure depends strongly on the size of the A-site and B-site ions. If there is a size mismatch between the A-site and B-site ions and the space in the lattice where they reside, the perovskite structure will become distorted. Such a lattice distortion of perovskite structure is governed by Goldschmidt's tolerance factor (τ) (Goldschmidt, 1953).

$$\mathbf{r} \equiv (\mathbf{r}\_{\mathrm{A}} + \mathbf{r}\_{\mathrm{O}}) / \sqrt{2} (\mathbf{r}\_{\mathrm{B}} + \mathbf{r}\_{\mathrm{O}}) \tag{1}$$

where rA and rB are the mean radii of the ions occupying the A-sites and B-sites, respectively, and rO is the ionic radius of oxygen. For an ideal cubic perovskite *τ*=1.

Fig. 2. Unit cell structure of prototype cubic perovskite, ABO3, with A ions situated at the corners, B ions in the center and oxygen ions at the centers of the faces.

If *τ* differs slightly from unity the atoms are displaced from their ideal positions to minimize the free energy and a distorted perovskite structure is formed. For 0.96<*τ*<1 the lattice structure transforms to the rhombohedral structure and then to orthorhombic structure for *τ*<0.96, in which B-O-B bond angle gets distorted from 180. Structure like tetragonal, hexagonal and monoclinic etc. are reported for manganites for different x values. Among the studied samples the LSMO, LBMO and LNMO systems crystallize in the rhombohedral with *R3c* space group (Venkataiah et al., 2007a, 2010; Kalyana Lakshmi and Reddy 2009a) while LCMO and NSMO group of samples crystallize in to orthorhombic structure with *Pbnm* space group (Venkataiah et al., 2005; Venkataiah and Reddy 2009b). A typical XRD patterns of LSMO & NSMO samples are shown in Figure 3. The X-ray linewidths are used to estimate average crystallite size <s> values through the Scherrer's formula modied by Williamson and Hall (1953).

octahedral interstitial sites. The BO6 octahedra make contact to each other by their vertices and form a three-dimensional network. In manganites the A-sites of the perovskite structure are occupied by trivalent rare earth ion while B-sites with Mn ions. The stability of the perovskite structure depends strongly on the size of the A-site and B-site ions. If there is a size mismatch between the A-site and B-site ions and the space in the lattice where they reside, the perovskite structure will become distorted. Such a lattice distortion of perovskite

where rA and rB are the mean radii of the ions occupying the A-sites and B-sites,

Fig. 2. Unit cell structure of prototype cubic perovskite, ABO3, with A ions situated at the

If *τ* differs slightly from unity the atoms are displaced from their ideal positions to minimize the free energy and a distorted perovskite structure is formed. For 0.96<*τ*<1 the lattice structure transforms to the rhombohedral structure and then to orthorhombic structure for *τ*<0.96, in which B-O-B bond angle gets distorted from 180. Structure like tetragonal, hexagonal and monoclinic etc. are reported for manganites for different x values. Among the studied samples the LSMO, LBMO and LNMO systems crystallize in the rhombohedral

while LCMO and NSMO group of samples crystallize in to orthorhombic structure with *Pbnm* space group (Venkataiah et al., 2005; Venkataiah and Reddy 2009b). A typical XRD patterns of LSMO & NSMO samples are shown in Figure 3. The X-ray linewidths are used to estimate average crystallite size <s> values through the Scherrer's formula modied by

*3c* space group (Venkataiah et al., 2007a, 2010; Kalyana Lakshmi and Reddy 2009a)

corners, B ions in the center and oxygen ions at the centers of the faces.

with *R*

Williamson and Hall (1953).

τ= (rA+rO)/2(rB+rO) (1)

structure is governed by Goldschmidt's tolerance factor (τ) (Goldschmidt, 1953).

respectively, and rO is the ionic radius of oxygen. For an ideal cubic perovskite *τ*=1.

Fig. 3. (a, b) XRD patterns of LSMO & NSMO group of manganites at various sintering temperatures (TS) (c) Williamson-Hall plot of Nd0.67Sr0.33MnO3 manganites.

In this method, the experimental x-ray peak broadening is given by =tan +K/ st cos, where is the full width at half maximum (FWHM) of the XRD peaks after subtracting FWHM of the standard sample (SiO2), is the Bragg angle, K is the grain shape factor (0.89), is the wavelength of Cu K radiation, and is the strain in the sample. A plot between cos versus sin is shown in Figure 3(c). The average crystallite sizes were calculated from the intercept of the straight line with the vertical axis, while the slope of the line gives the lattice strain (). The average crystallite sizes are found to increase with increasing sintering temperature and are in nanometer range. The variation of crystallite size with sintering temperature is shown in Figure 4. In fact, the computed crystallite sizes are comparable with those obtained from SEM. Figure 5 shows the representative SEM images of LSMO samples sintered at 800 and 1100ºC. The SEM observation reveals that there is a uniform distribution of grain sizes for the samples and as the sintering temperature increases, the grain size increases and porosity decreases.

Influence of Sintering Temperature

magnetotransport properties and microstructure.

Fig. 6. Magnetic and electronic phase diagram versus doping x for La1-xCaxMnO3

manganites (Schiffer et al., 1995, Urushibara et al., 1995).

on Magnetotransport Behavior of Some Nanocrystalline Manganites 273

structural, magnetic and electronic phase transitions occur simultaneously. The magnetic and electronic properties of lanthanum manganites, La1-xAexMnO3 (Ae = Ca, Sr, etc) are strongly related to the simultaneous presence of manganese in different valence states. The end-members, *x* = 0 and *x* = 1, containing Mn in only one valence state (for divalent ion at Asite substitution leads to Mn in 3+ and 4+ valence states), are usually antiferromagnetic insulators (AFI), but intermediate compositions, which have mixed Mn valence, may be ferromagnetic and have good conducting properties. Figure 6 illustrates the magnetic and electrical properties of La1-xCaxMnO3 (x=0-1) system with change in composition and temperature (Schiffer et al., 1995, Urushibara et al., 1995). Ferromagnetic behavior starts to manifest itself when x is 0.1 and the composition upto x~0.3 have both antiferromagnetic (AFM) and ferromagnetic (FM) characteristics. The x=0.3 composition is clearly FM, while the x>0.5 compositions are AFM. The coexistence of metallic conductivity and ferromagnetic coupling in these materials at low temperature has been explained in terms of a double exchange mechanism, proposed by Zener in 1951. The double exchange mechanism involves simultaneous transfer of an electron from Mn3+ to O2- and from O2- to Mn4+. Moreover, the properties of manganites are not only sensitive to the manganese valency, but also are strongly affected by other factors such as average cationic radius <rA> of A-site and A-site cationic mismatch quantified by σ2 which effects the magnetic interaction. Due to the long history of work on these compounds, most of the studies have been performed on bulk ceramics and thin films. Apart from this it was demonstrated that the properties of manganites are strongly affected by particle size. It was found that the sintering temperature is one of the key factors which influences the crystallization and microstructure of the perovskite samples. The basic magnetic properties such as spontaneous magnetization, the Curie temperature and coercivity are strongly influenced with sintering temperature. Mahesh et al., (1996) were the first to explore the size effects in polycrystalline La0.67Ca0.33MnO3prepared by citrate-gel route and observed a close relationship between

Fig. 4. Variation of Crystallite size (<s>) in nanometers with Sintering temperature (TS C).

#### **3.2 Magnetic properties**