Powder Metallurgy Progress, Vol.4 (2004), No 1 57 INTERMETALLIC ALLOYS BASED ON Ni-Al PRODUCED BY POWDER METALLURGY E. Bobrová, M. Besterci Abstract The paper summarises knowledge in the field of intermetallic alloys produced by powder metallurgy (PM). It presents some intermetallic alloys, their classification and practical use. It deals with the technology of production of intermetallic alloys by PM, such as reactive sintering, CIP, reactive HIP, hot extrusion pressing, injection moulding, etc. Pre- alloyed powders or mixtures of elemental powders with potential admixtures of dispersed or composite particles intended to strengthen the alloys were used as a starting material for the production of intermetallic alloys. The aim of the study was to prepare intermetallic alloys based on Ni-Al using reactive sintering or HIP. The produced intermetallic alloys Ni 3 Al, Ni 3 Al+1wt. % Al 2 O 3 , Ni 3 Al + 2 wt. % TiB 2 were investigated for their mechanic properties in relation to microstructure and mechanism fracture. Keywords: intermetallic alloys, powder metallurgy, reactive sintering, CIP, HIP, hot extrusion pressing, injection moulding INTRODUCTION Due to their attractive properties, intermetallic alloys constitute a group of modern and progressive materials. The use of superalloys with a good combination of strength and toughness is limited to temperatures below 1100°C. The modern engineering ceramics can be used at much higher temperatures, but they are inherently brittle because of the covalent atomic bonding. The intermetallics are expected to fill the gap between the superalloys and the ceramics, not only with respect to the service temperature range, but also with respect to the properties; higher strength than the superalloys because of the strong atomic bonding, and less brittleness than the ceramics because the bonding is still metallic [1]. Intermetallic alloys are new-generation construction materials resistant to high temperatures and oxidation and known for very interesting physical and mechanical properties. Presently they are subject to intensive research in all developed countries a round the world. This research is concentrated in universities and specialised research centres particularly in the USA, Canada, Japan, Germany, France and Switzerland. The research of these intermetallic alloys is supported particularly by aviation and automobile industry companies [2]. Up to this date, a range of intermetallic alloys based on Fe, Al, Ti, Nb, La, Zr, Li, Co, Sn, Au, Ga, Ge, Pd and Hf has been developed. These alloys are either simple, most frequently binary or ternary, or composite, strengthened with particles or fibres [3]. Intermetallic phases had been in use for various purposes for many centuries because of their comparatively high hardness, Tab.1, whereas in modern times their particular physical properties have been of primary interest, Tab.2 [4,5]. Eva Bobrová, Michal Besterci, Institute of Materials Research of SAS, Košice, Slovak Republic
Transcript
Powder Metallurgy Progress, Vol.4 (2004), No 1 57
INTERMETALLIC ALLOYS BASED ON Ni-Al PRODUCED BY POWDER METALLURGY
E. Bobrová, M. Besterci
Abstract The paper summarises knowledge in the field of intermetallic alloys produced by powder metallurgy (PM). It presents some intermetallic alloys, their classification and practical use. It deals with the technology of production of intermetallic alloys by PM, such as reactive sintering, CIP, reactive HIP, hot extrusion pressing, injection moulding, etc. Pre-alloyed powders or mixtures of elemental powders with potential admixtures of dispersed or composite particles intended to strengthen the alloys were used as a starting material for the production of intermetallic alloys. The aim of the study was to prepare intermetallic alloys based on Ni-Al using reactive sintering or HIP. The produced intermetallic alloys Ni3Al, Ni3Al+1wt. % Al2O3, Ni3Al + 2 wt. % TiB2 were investigated for their mechanic properties in relation to microstructure and mechanism fracture. Keywords: intermetallic alloys, powder metallurgy, reactive sintering, CIP, HIP, hot extrusion pressing, injection moulding
INTRODUCTION Due to their attractive properties, intermetallic alloys constitute a group of modern
and progressive materials. The use of superalloys with a good combination of strength and toughness is limited to temperatures below 1100°C. The modern engineering ceramics can be used at much higher temperatures, but they are inherently brittle because of the covalent atomic bonding. The intermetallics are expected to fill the gap between the superalloys and the ceramics, not only with respect to the service temperature range, but also with respect to the properties; higher strength than the superalloys because of the strong atomic bonding, and less brittleness than the ceramics because the bonding is still metallic [1].
Intermetallic alloys are new-generation construction materials resistant to high temperatures and oxidation and known for very interesting physical and mechanical properties. Presently they are subject to intensive research in all developed countries a round the world. This research is concentrated in universities and specialised research centres particularly in the USA, Canada, Japan, Germany, France and Switzerland. The research of these intermetallic alloys is supported particularly by aviation and automobile industry companies [2]. Up to this date, a range of intermetallic alloys based on Fe, Al, Ti, Nb, La, Zr, Li, Co, Sn, Au, Ga, Ge, Pd and Hf has been developed. These alloys are either simple, most frequently binary or ternary, or composite, strengthened with particles or fibres [3].
Intermetallic phases had been in use for various purposes for many centuries because of their comparatively high hardness, Tab.1, whereas in modern times their particular physical properties have been of primary interest, Tab.2 [4,5].
Eva Bobrová, Michal Besterci, Institute of Materials Research of SAS, Košice, Slovak Republic
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The high hardness is correlated with high strength at high temperatures, and thus intermetallic phases have been regarded as promising with respect to the development of new structural materials for high-temperature applications [1].
Tab.1. Some applications of intermetallic phases in past centuries.
Phase Application Cu3As Coating of bronze (Egypt) Cu31Sn8 Mirrors (China) Sn8Hg Mirrors surface (Italy) Ag2Hg3+Sn6Hg Dental amalgam (China) Cu4Hg3 Dental amalgam (Germany) SbSn Type metal
Tab.2. Some applications of intermetallic phases.
Phase Application since about Ni3Fe FeCo(-2V) Fe3(Si,Al) SmCo5
High permeability magnetic alloy (Permalloy) Soft magnetic alloy (Permendur) Magnetic head material (Sendust) Permanent magnets
1920 1930 1935 1970
Nb3Sn A 15 superconductors 1965 CuZnAl CuNiAl NiTi
Shape-memory alloys Shape-memory alloys (Nitinol)
1960 1965
MoSi2 Heater elements (Mosilit, Super Kanthal)
1955
NiAl, CoAl Protective coatings 1965
CLASIFICATION OF HIGH-TEMPERATURE INTERMETALLICS An intermetallic compound is generally an ordered alloy phase formed between
two metallic elements. Intermetallic compounds are generally composed of two elements on ordered sites in a variable proportion around a few stoichiometric ratios. Presently, the majority of intermetallic compounds can be classified as being formed by two elements, A and B, into five different stoichiometric combinations with compositions around A3B, A2B, A5BB3, A7B6B and AB. Within each stochiometric group, the intermetallic compounds can occur with a different crystal structure type. Some of the commonly occurring crystal structures among high-temperature intermetallic compounds are listed in Table 3 [6-8].
Intermetallics have recently received considerable attention as potential structural materials for many engineering applications. They have a number of properties, which make them very attractive for high temperature applications [9].
These properties are as follows: • Intermetallics are intrinsically strong and the strength does not degrade rapidly with
increasing temperature. Compounds having the L12 (ordered f.c.c.) and some with B2 (ordered b.c.c.) crystallographic structures show an increase in yield and flow stresses with increasing temperature.
• Those intermetallics formed from various combinations of aluminium, titanium and silicon elements, i.e. aluminides, titanides and silicides, have low and very low
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densities. This gives rise to high specific properties (modulus and strength divided by density) which are particularly important for aerospace and transportation applications.
• Elastic moduli of intermetallics are quite high and they tend to decrease more slowly with increasing temperature than those of disordered alloys.
• Intermetallics with high aluminum content are expected to have good oxidation resistance at high temperatures.
• Due to their ordered structures, intermetallics exhibit lower self-diffusion coefficients and these result in a slower rate of creep, recrystallization and corrosion.
Tab.3. Classification of high-temperature intermetallic compounds.
Stoichiometry Crystal Structure Example Melting Point [oC]
Density [g/cm3]
A3B
L12
DO19DO22
A15
A12
N3Al Pt3Al Ti Sn 3Ni3Ta Al Nb 3Al3Ta Nb Al 3Mo Si 3V Si 3Cr3Si
The single largest disadvantage of intermetallics is their low ductility, particulary
at ambient to intermediate temperatures. Factors, which may lead to a limited plasticity (ductility) of these materials, and eventually brittle fracture, are numerous. The most fundamental factors include a limited number of easy slip systems, a large slip vector, restricted cross-slip, difficulty of transmitting slip across grain boundaries, segregation of
Powder Metallurgy Progress, Vol.4 (2004), No 1 60 harmful impurities to grain boundaries, intrinsic weakness of grain boundaries due to their structural configurations, and finally, environmental effects [10-18].
POWDER METALLURGY OF INTERMETALLIC ALLOYS One of the most progressive ways of producing intermetallic alloys with a
complex of high-quality properties is powder metallurgy (PM). The products obtained by powder metallurgy have a range of advantages, for example, from an economical point of view, their production is less demanding compared to conventional metallurgy. The most important advantages of this technology include the following: lower energy consumption, lower workability, higher production, higher material homogeneity and a lower production of scrap. The PM technology allows one to prepare materials with fine and homogeneous microstructure, which cannot be obtained by conventional melt metallurgy. The conventional technological approaches are frequently unable to keep up with the intensive industrial development and the basic strengthening methods, such as plastic deformation, alloying, and precipitation, appear insufficient to comply with the demanding conditions to which the parts are exposed in practical life. The strengthening methods are effective up to the temperatures 0.3 - 0.4 Tmelt. However, when the materials are to be used at temperatures above 0.5 Tmelt, it is necessary to strengthen them with fibres or dispersed particles.
In addition to geometrical factors, the basic criteria for the strengthening phase is its stability in contact with the matrix. These properties are exhibited by difficult-to-melt compounds, such as oxides, carbides, nitrides and borides. Their melting point should exceed that of the matrix. The effectiveness of strengthening depends on the size of the strengthening particles (to not exceed 10-50 nm), medium distance between strengthening particles (100-500 nm), and their volume proportion (1-10 %). The use of dispersed particles strengthens the matrix, but on the other hand, decreases its plastic properties. Because of that, it is necessary to select such volume proportions of the strengthening phase, which ensure a good combination of high-temperature strengthening and refractoriness [19,20].
The starting materials used for production of intermetallic alloys include elemental powders and pre-alloyed powders. In order to increase their physical and mechanical properties, these powders may be strengthened with dispersed particles or fibres. In the preparation of composite materials, dispersed particles such as TiBB2, Al2O3 and Y2O3 or fibres such as SiC or Al2O3 have been used. They are difficult-to-melt compounds with high melting points. In general, dispersed oxides have some advantages compared to borides, carbides or nitrides because a wide range of dispersity can be achieved with oxides and, at the same time, their introduction to the matrix is simpler due to their chemical stability towards the matrix. The mixtures prepared are homogenised in a vibrating mill and consolidated by several technological methods, such as reactive sintering, CIP, reactive HIP, hot pressing, explosive compacting, hot extrusion pressing, rolling, and injection moulding. The final operation used in many cases is the thermal processing intended to increase the homogeneity of the structure and to improve the mechanical properties. The subsequent sections deal with some selected technologies of the preparation of intermetallic alloys [21].
Reactive sintering One of the most interesting techniques utilised to consolidate elemental powders is
reactive sintering. Reactive sintering is a technological procedure used for the production of intermetallic alloys, particularly those based on nickel-aluminium. During this procedure, a transient liquid phase develops by heating the compressed mixture of elemental powders.
Powder Metallurgy Progress, Vol.4 (2004), No 1 61 This liquid phase differs from that during conventional sintering because it vanishes before the onset of cooling, and the stage of solidification. In this way, an intercrystalline compound is produced and the compact is sintered at the same time. In the solid state, the process of sintering is controlled by diffusion, which is relatively slow, even at high temperatures, and depends on individual atomic movements. This occurs in pure metals, powder mixtures prepared from elemental or pre-alloyed powders which are homogenised more rapidly. Contrary to that, the diffusion in the liquid state occurs at a much faster rate than in the solid state [21].
Average particle sizes of Al and Ni were 3 and 15 μm, respectively. Exposures of stoichiometric compacts to a furnace temperature between 500 and 800°C provide material with ∼96 % density in a few minutes. An exothermic reaction that occurs during sintering rapidly raises metal temperature to ∼1500°C. The pores can be completely removed by subsequent HIP runs; however, the compacts also display about 5 % of a second phase, probably Ni5Al3. Much of this phase can be removed by annealing in the range 1000-1350°C [22].
Figure 1 shows a schematic binary phase diagram for a reactive sintering system, where a stoichiometric mixture of A and B powders is used to form an intermediate compound product AB. The reaction occurs above the lowest eutectic temperature in the system, yet at a temperature where the compound is solid. At the lowest eutectict temperature a transient liquid forms and spreads through the compact during heating. Generally, heat is liberated because of the thermodynamic stability of the high melting temperature compound. Consequently, reactive sintering is nearly spontaneous once the liquid forms [23].
Fig.1. A schematic binary phase diagram showing the characteristics necessary for formation of an intermetallic compound AB from mixture A and B constituent powders.
Fig.2. A sketch of the envisioned reactive sintering process for forming Ni3Al from mixed elemental Ni and Al powders. As temperature increases, first a solid state reaction occurs, with subsequently a rapid reaction when the eutectic liquid forms. The final product is a densified compound [22].
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Figure 2 shows a schematic diagram of the reactive sintering process as envisioned from this research [22]. Nickel and aluminium powders are randomly mixed in a stochiometric ratio. The powders have small particle sizes to aid intermixing and are milled and compressed to create good particle-particle contact. This mixture is sintered under precise conditions of atmosphere, heating rate, time and temperature. During heating to the first eutectic temperature, solid state interdiffusion generates intermetallic compound phases and some self-heating. At the first eutectic temperature, liquid forms and rapidly spreads throughout the structure. The liquid consumes the elemental powders and forms a precipitated Ni3Al solid behind the advancing liquid interface. Interdiffusion of nickel and aluminium is quite rapid in the liquid phase, and the compound generates heat, which further accelerates the reaction. If the reaction is controlled, then the compound will be nearly fully densified and suitable for containerless hot isostatic compaction to full density. Success with this process is measured by a low final porosity, good shape retention, and good mechanical properties. Figure 3 provides a schematic flow chart of this fabrication process, listing the key steps and variables [23-28].
Fig.3. A schematic diagram of the fabrication process used for reactive sintering Ni3Al [23].
Fig.4. Schematic diagram of the injection-molding process [25].
Injection moulding Reactive consolidation also may be applied to injection moulding, a process
increasingly being used for the powder processing of both monolithic and composite alloys. In injection moulding, a mixture of powders, short fibres and a binder is extruded through a tapered die to achieve fibre alignment, Fig.4. Extrusion must be performed above the softening temperature of the binder. After extrusion the binder is removed (thermally or by
Powder Metallurgy Progress, Vol.4 (2004), No 1 63 wicking action) and the compact is consolidated to approximately full density by HIP. Apart from the alignment of fibres, which is achieved only when particles and fibres are very small in diameter (10 μm), this process offers the possibility of producing complex powder metallurgy parts. However, the principal disadvantages are the difficulty of complete binder removal and the inability to produce continuous-fibre-reinforced composites. Nevertheless, fully dense Al2O3 – reinforced composites of NiAl have been produced successfully by this method. The latter composite was made from prealloyed powders. There was some evidence that the Al2O3 fibres inhibited crack propagation in the NiAl matrix [25,26,29].
HIP, HIP + extrusion Powders can be consolidated either by hot pressing, HIP, HIP + extrusion, or by
direct extrusion of powders. Hipping is usually carried out at 1100-1150°C for 3 h, with a pressure of 120 MPa. Hipping leads to the 100 % density with grain size ∼10-15 μm for ternary and quaternary alloys. Direct extrusion with a reduction ratio of 8:1 results in 100 % density. Billets up to a 200 mm diameter, compacted under atmospheric pressure (CAP) generally contain large non-interconnected porosity, but can be further consolidated by extrusion at 1100°C in a stainless steel can. Complex shapes made by CAP can be further consolidated by hipping. Wright and Flinn reported that extruded powder generally has higher strength and ductility than the HIP material [26].
MATERIALS AND EXPERIMENTS The experiments were carried out on specimens produced in two different ways,
conventional pressing and sintering, and reaction sintering with isostatic pressing and sintering. The study [30] focused on preparation of the intermetallic alloy Ni3Al by PM, also
using consolidation by reactive sintering. The starting material consisted of elemental powders at the following ratio: 86.72 wt. % Ni : 13.28 wt. % Al. We used < 5 μm fraction of Ni-powder and two fractions, < 40 μm and < 20 μm, of Al powder. The nickel and aluminium powders were used to prepare two variants of powder mixtures, which were homogenised in a vibration mill for 30 min. The specimens intended for static tensile tests were compacted at pressures of 200, 250, 300 and 400 MPa, and the sintering temperatures were 400, 450, 470, 500, 570, 600, 650°C, applied for 30 and 60 min at a heating rate of 20°C /min. The sintering was carried out in a vacuum of 10-6 Pa.
Hot isostatic pressing produced three types of Ni3Al intermetallic alloys. The first type was the basic material containing only Ni and Al. It was prepared from nickel powder (40 μm mesh) and aluminium powder (20 μm mesh) at a ratio of 86.72:13.28 wt. % in order to obtain the Ni3Al intermetallic alloy. The elemental powders were mixed in an attritor for 30 min. in order to homogenise the mixture. The homogenised mixture was sealed in steel capsules with the dimensions of 20 mm dia x 250 mm after evacuation to a pressure of 10-5 Pa. The second type had a modified composition obtained by adding 1 wt. % of pure Al2O3 powder to the basic material. The third type was produced by the addition of 2 wt. % of TiB2 powder. The processing of both the second and third types was identical to the processing of the basic material. The melting point of the Al2O3 powder was 2045°C and the density reached 3. 95.103 kg/m3. The TiBB2 powder has a melting point of 2900°C and a density of 4.38.10 kg/m . All three types of the material were compacted by hot isostatic pressing at a pressure of 150 MPa and sintered at 1150°C for 1 hour [1,2]. After removing the steel sheath from the compacted material, specimens for tensile testing with dimensions of 4 mm dia x 20 mm gauge length were machined.
3 3
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A Neophot optical microscope was used for microstructural analysis. The quantitative analysis was carried out on a TESLA 340 scanning electron microscope in order to determine the distribution of the elements in the phases using a LINK ISIS 300 microanalyser. Vickers microhardness HV 0.5 was measured on the HIPed samples [31-33], and the Vickers microhardness HV 0.01 was measured on the only sintered specimens.
RESULTS AND DISCUSSION The development of total porosity was studied in specimens produced by
conventional pressing and sintering in dependence on compacting pressure, sintering temperature and the sintering time. The mechanical mixture prepared with the <40 μm fraction of Al-powder was compacted at 200 to 400 MPa, and then sintered at 400°C for 30 min in a vacuum 10-6 Pa. Metallographical examination showed that the pores occurred as aggregates and formed irregular, chain-like formations. The total porosity reached rather high levels ranging between 28 and 30 %. A decrease in total porosity to 18-20 % was observed in the prepared specimens (preparation condition 400 MPa, 450°C for 30 min in a vacuum 10-6 Pa) and a change in the shape of pores was also observed. The pores appeared as smaller irregular formations. Neither in the first nor in the second case did complete transformation to Ni3Al occur due to low temperatures of sintering, a short period of sintering, or the use of a coarser fraction of Al-powder.
The mechanical mixture prepared with the < 20 μm fraction of Al-powder was also compacted and sintered (200 to 400 MPa, 400 to 650°C for 60 min in a vacuum 10-6 Pa). The results obtained allowed us to state that reactive sintering occurred during the process of consolidation, and the optimum preparation parameters for the respective mechanical mixture were the following: compaction at 400 MPa, sintering at 650°C for 60 min in a vacuum of 10-
6 Pa. The resulting microstructure consisted of polyedric grains γ´ (Ni3Al) of size 15-18 μm, Fig.5. At this temperature we can already talk about the complete spheroidisation of pores. The total porosity declined down to 4 to 5 %. The microhardness reached the values of 394-700 HV0.01. The values of strength and plastic characteristics are presented in Tab.4. The fracture surfaces obtained by static tensile tests showed intercrystalline separation, Fig.6, which was manifested by the low level of plastic properties.
Tab.4. Strength and plastic characteristics of the Ni3Al material (400 MPa, 650°C for 60 min in a vacuum of 10-6 Pa, at room testing temperature).
Material Ultimate tensile strength Rm [MPa]
Yield strength Rp0.2 [MPa]
Ductility A5 [%]
Microhardness HV 0.01
Ni3Al 310-450 250-280 3-4 394-700
All three types of material produced by hot isostatic pressing had high density and thus low residual porosity. The measured porosity was below 1 %. The microstructure of the Ni3Al material was relatively homogeneous with fine polyedric grains of size 5-10 μm, Fig.7. The few inhomogeneities observed by optical microscopy, subsequently analysed by the LINK microanalyser, showed to be Ni-rich areas, Fig.8.
The origin of these microstructural inhomogeneities is the incomplete transformation of the constituents Ni and Al into the Ni3Al intermetallic as a result of insufficient temperature and short compaction time [3]. The microhardness measurements reflected the microstructural inhomogeneity: in the homogeneous areas, the values of 500-800 HV 0.5 were registered, but in the inhomogeneous ones, the microhardness increased to 1500-1600 HV 0.5. Table 5 summarizes the values of HV 0.5 for both the homogeneous
Powder Metallurgy Progress, Vol.4 (2004), No 1 65 areas without pronounced inhomogeneities, and for the areas of microstructural changes caused by chemical inhomogeneity. The results presented in Tab.5 show that in the materials Ni3Al + 1 wt. % Al2O3 and Ni3Al + 2 wt. % TiB2, the microhardness increased to 700-950 HV 0.5 and 840-1050 HV 0.5, respectively. This increase was reflected also in the areas of inhomogeneous microstructure, but less markedly, Tab.5.
Fig.6. Fracture surface of the sample Ni3Al after compaction at 400 MPa and sintering at
650°C.
In the material Ni3Al + 1 wt. % Al2O3, fine polyedric homogeneous microstructure with a grain size below 20 μm was also observed. The specific feature of this microstructure was the homogeneously distributed dispersoids along grain boundaries. In the material Ni3Al + 2 wt. % TiB2, the grain size was in the range between 10-20 μm, but the dispersoid particles were not homogeneously distributed along the grain boundaries and formed clusters in certain areas [4].
Strength and plasticity were also analysed for the three investigated materials. The ultimate tensile stress (U.T.S) of the basic Ni3Al material was 419 MPa, measured at room temperature. The ductility A5 = 3 % was very low. The tensile test at 700°C revealed a decrease in U.T.S to 198 MPa, and in ductility to 1 %. The material Ni3Al + 1 wt. % Al2O3 exhibited an increase in tensile strength by one third compared to the basic material. The plasticity decreased as a result of the addition of disperses particles [5].
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Fig.7. Microstructureof the sample Ni3Al
produced by HIP. Fig.8. Inhomogeneities in the Ni3Al
material.
The material dispersion strengthened by TiBB2 also showed an increased strength. The measured values of tensile strength and ductility are listed in Tab.5, for both room temperature and 700°C. The fractographical analysis of the fracture surfaces of the tensile specimens tested at room temperature showed two types of fracture mechanisms for the basic Ni3Al material. Intercrystalline separation dominated 60 % of the total fracture surface, Fig.9, while the rest of the fracture surface showed transcrystalline cleavage. The fracture surfaces of the specimens tested at 700°C exhibited intercrystalline separation as a basic mechanism, Fig.10.
Fig.9. Fracture surface of the Ni3Al material. Fig.10. Intercrystalline separatio in the Ni3Al.
The fracture surface of the tensile specimens prepared from the material Ni3Al + 1 wt. % Al2O3, and tested at room temperature, showed that the basic fracture mechanism was intercrystalline separation with a very restricted presence of transcrystalline cleavage. The fracture formed at 700°C showed only intercrystalline separation.
The fracture surface of the material Ni3Al + 2 wt. % TiB2 tested at room temperature and at 700°C was equally intercrystalline. In certain areas of the fracture surface, inhomogeneously distributed clusters of the TiB2 dispersoid were observed.
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CONCLUSIONS In conclusion we can state that an intermetallic alloy Ni3Al was produced by PM
technology using reactive sintering. The most suitable fractions of powders from among the available ones were Ni-powder, the fraction < 5 μm, and Al-powder of a fraction < 20 μm, and the optimum parameters of compacting were as follows: compaction pressure of 400 MPa, sintering at 650°C for 60 min in a vacuum of 10-6 Pa. Under these conditions, a Ni3Al intermetallic alloy was produced with total porosity reaching 4 to 5 %. The strength properties reached the following values: Rm = 450 MPa, Rp0.2 = 280 MPa and ductility A5 = 4 %. The microstructure was composed of fine polyedric grains of size 15-18 μm and the fracture surface exhibited intercrystalline separation.
The materials Ni3Al, Ni3Al + 1 wt. % Al2O3 and Ni3Al + 2 wt. % TiB2 had a fine and relatively homogeneous microstructure. The residual porosity was below 1 %. The alloys exhibited relatively high values of microhardness. The strength properties of the dispersion strengthened alloys were by one third higher than those of the basic intermetallic alloy Ni3Al. On the other hand, the plasticity characteristics were very low in all cases. A common feature of the three alloys is their high brittleness and low plasticity due to intercrystalline damage. The reason of this elevated brittleness is the high resistance to yielding, which can be eliminated with B addition and increases the cohesion strength of the grain boundaries and promotes slip nucleation. The segregation of disperse particles along grain boundaries contributes to the elevated brittleness.
Acknowledgements The authors are grateful to the Slovak Grant Agency for Science (VEGA
Germany, September 7-11, 2003, p.147 [30] Bobrová, E., Besterci, M.: AMS, vol. 1, 1995, p.122 [31] Bobrová, E., Besterci, M.: Journal of Mining and Metallurgy B, vol. 36 , 2000, p.55 [32] Bobrová, E., Besterci, M., Sülleiová, K. In: Deformation and Fracture in Structural PM
