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Reaction sintering of NiAl and TiB2–NiAl composites under pressure
Materials Science and Engineering A257 (1998) 341 – 348
Reaction sintering of NiAl and TiB2 –NiAl composites under pressure S.K. Bhaumik, C. Divakar, L. Rangaraj, A.K. Singh * High Pressure Laboratory, Materials Science Di6ision, National Aerospace Laboratories, Bangalore 560 017, India Received 7 May 1998
Abstract Intermetallic matrix composites are a new class of engineering materials for high temperature structural applications in oxidizing and aggressive environments. Attempts have been made in the present investigation to synthesize TiB2 –NiAl composites, in which the matrix phase NiAl was produced in situ by reaction synthesis. The composites with 10, 15 and 30 vol.%NiAl were fabricated from the mixtures of elemental Ni and Al and TiB2 powders by high pressure reaction sintering (HPRS) and reactive hot pressing (RHP). The HPRS and RHP were carried out at 3 GPa and 900°C, and 20 MPa and 1650°C, respectively. The different phases in the sintered compacts were identified by X-ray diffraction and microstructural studies. In HPRS, the reaction was incomplete which gave rise to various intermediate phases (Ni2Al3, Ni3Al). It was necessary to anneal these compacts at 1100°C to obtain TiB2 –NiAl composites with single phase NiAl matrix. The densities of the HPRS compacts were 99%. The hardness and fracture toughness values were in the range 10 to 20 GPa and 3.9 to 5.7 MPa m, respectively. The RHP compacts contained AlB2, Ni2B and NiTi2 phases in addition to those present in the HPRS compacts. The RHP composites were fully dense. These had superior hardness (15 – 22 GPa) but inferior fracture toughness (2.9 – 3.8 MPa m) compared to those obtained by HPRS. © 1998 Elsevier Science S.A. All rights reserved. Keywords: Composites; Sintering; Matrix phase; High pressure
1. Introduction The unique combination of high melting point (2900°C), high hardness (27 GPa), high elastic modulus (370 GPa), high electrical conductivity, and relatively low coefficient of thermal expansion (4.6× 10 − 6°C) has made TiB2 an important material for high performance applications [1–4]. Attempts have been made [5–7] to consolidate TiB2 powder into dense solid bodies by reaction sintering, and hot pressing with or without sintering aids. In most cases, the consolidation by solid state sintering requires temperatures of 2100°C and times over 60 min. This often results in exaggerated grain growth leading to inferior mechanical properties. The sintering temperature and time can be reduced with the use of liquid phase forming additives. In such cases, * Corresponding author: Tel.: +91 80 5270098; fax: + 91 80 5260826; e-mail: [email protected]
a low melting phase is formed at the grain boundaries which adversely affects the high temperature mechanical properties of the products. Nickel aluminide containing more than 41 at.%Al forms a single phase ordered B2 structure based on the body centered cubic (bcc) lattice [8]. In terms of thermophysical properties, NiAl (B2) is more suitable for high temperature applications than Ni3Al (L12) [8,9]. It has high melting point (1638°C), low density (5.86 g cm − 3), high young’s modulus (294 GPa), and high thermal conductivity (76 Wm − 1 K − 1) at ambient temperature [10–12]. In addition, NiAl has excellent oxidation resistance at high temperatures (1000°C). However, the structural use of NiAl suffers from poor fracture resistance at ambient temperatures, and low strength and creep resistance at elevated temperatures. Various processing techniques such as rapid solidification [13], mechanical alloying [14], reaction sintering [11,15] and RHP are normally used for the synthesis of
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NiAl intermetallic. HPRS of stoichiometric and Ni-rich NiAl has also been reported recently by Cheng et al. [16]. The Ni–49.3 at.%Al produced by HPRS at 2 GPa and 500°C is reported [16] to have high density (98.2%)), good compressive ductility (14.5%), and strength (876 MPa). Because of the good high temperature properties of NiAl, efforts have been made to use this intermetallic as a matrix for high temperature composites. Various composites of NiAl with Al2O3, Y2O3 and TiB2 have been reported [11,17 – 20]. The effects of fine TiB2 (1–3 mm) particles on NiAl have been studied in detail [17–20]. In these studies, the addition of the refractory particulate was limited to maximum 30 vol.%. These particles improve the compressive strength of stoichiometric NiAl significantly. In the present investigation,TiB2 – NiAl (10–30 vol.%) composites have been prepared by HPRS and RHP. The stoichiometric composition of NiAl has been chosen as the matrix phase. The results are presented in this paper.
The RHP was conducted at 1650°C in flowing argon atmosphere for 1–30 min. The heating was done by an induction furnace. The heating rate was 15°C s − 1. After HPRS and RHP, the reaction products were identified by X-ray diffraction. The density was measured by volume displacement method. The theoretical density used for comparison was calculated by the volumetric rule of mixtures assuming that Ni and Al reacted completely to form NiAl. The microstructural observations were carried out in optical (Neophot-2) and JEOL 840A scanning electron microscopes (SEM). Shimadzu HSV-20 microhardness tester with a load of 1.96 N was used to determine the hardness. For indentation fracture toughness determination, a load of 49 N was used. The fracture toughness was calculated from the equation [23]: KIC = 0.0824 (P/C 3/2) where, P is the indentation load (N) and C the crack length (mm). Ten indentations were made on each sample for hardness and fracture toughness measurements, and the average values are reported.
2. Experimental 3. Results and discussion The raw materials used in the present study were elemental nickel and aluminium, and TiB2 powders. A master alloy powder mixture of Ni – 50 at.%Al was first prepared by ball milling (Fritsch Pulverisette-5 centrifugal mill) in hexane medium for 24 h. This master alloy powder was used to prepare the composite powder mixtures of desired compositions. In the second step, the milling was carried out for 8 h. The details of the powder characteristics are given in Table 1. The powder mixture compositions and the synthesis conditions are given in Table 2. The HPRS of NiAl and TiB2 – NiAl composites was carried out in a 200 t cubic anvil apparatus (anvil face, 10 mm square) capable of generating pressures up to 6.5 GPa and temperatures up to 3000°C. The details of the high pressure-high temperature cells used in these experiments are reported elsewhere [21]. In an actual experiment, the sample was first pressurized to 3 GPa, and then heated to desired temperature (800 –900°C). The heating rate was 40°C s − 1. The temperatures were maintained for 15 – 45 min. The typical dimensions of the recovered samples were 3.5 mm in diameter and 4 mm in thickness. Some of these compacts were annealed at 1100°C in vacuum (10 − 4 Torr) for 120 min. The RHP of TiB2 – NiAl composites was carried out in a hot press fabricated in the laboratory [22]. The powder mixture was cold compacted at 400 MPa into a cylindrical compact of 13 mm diameter and 8 mm height. These compacts, with green densities in the range 55 to 60% were placed in a graphite die and a load corresponding to pressure of 20 MPa was applied.
3.1. Matrix phase NiAl (HPRS) The measured densities (as a percentage of theoretical density) of HPRS NiAl compacts fabricated at 800 and 900°C for 15 min are listed in Table 2. The maximum density obtained is 99.3%. During HPRS, the exothermic reaction between Ni and Al generates enough heat to produce a liquid phase. This liquid provides a faster diffusion than encountered in the solid state leading to enhanced densification. The amount of liquid phase formed depends on the powder composition (Al content) and temperature. It has been estimated [24] that stoichiometric Ni–Al composition can produce as much as 50 vol.% liquid when the temperature is in the range of 600–700°C. The application of pressure during reaction sintering helps in eliminating the big pores often observed in the pressureless sintering. It has been reported [25] that during the reaction synthesis of various intermetallics in the Ni–Al system, a pressure of 100 MPa is sufficient to obtain 99.9% dense compacts. In the present investigation, even though the pressure was much higher (3 GPa), the sintered compacts had 1% porosity. This porosity is due to the adsorbed gases and volatile impurities in the reactant powder mixtures. During synthesis, the released gaseous phases fill the pores and hinder complete densification. The high pressure cell is completely isolated from the atmosphere and the trapped gases can not escape from the compact. When the HPRS compacts are annealed at higher temperature (1100°C), the
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Table 1 Powder characteristicsa Material
TiB2
Nickel
Aluminium
Source
HC Starck, Germany
INCO (UK)
ALCOA
Chemical analysis
B (30.3) C (0.3) O (0.25) N (0.39) Ti (Balance)
C (0.06) Fe (0.005) O (0.05) Co (0.0003) N2 (0.03) S (0.003) Ni (Balance)
High purity atomized powder 7123
10
3.7
16.5
Particle size (mm) a
Al (99.3)
All compositions given in wt.%.
Table 2 Density, hardness and fracture toughness values of NiAl and TiB2–NiAl (10–30 vol.%) composites fabricated by HPRS and RHP Sample
Synthesis temperature (°C) time (min)
Density (% TD)
Hardness (Hv) (GPa) X (s)
X (s)
HPRS (NiAl) A B C (annealed)
800, 15 900, 15 900, 15
99.1 99.3 97.5
HPRS (TiB2–30 vol.%NiAl) D E F G (annealed)
900, 900, 900, 900,
15 30 40 30
99.2 99.1 99.1 98.7
10.5 13.2 14.2 12.8
(1.6) (0.6) (1.5) (0.6)
5.10 5.69 5.43 5.53
(0.60) (0.28) (0.41) (0.28)
HPRS (TiB2–15 vol.% NiAl) H I J K (annealed)
900, 900, 900, 900,
15 30 45 30
99.2 99.0 99.3 98.8
16.4 17.8 19.3 18.2
(1.4) (2.0) (1.1) (0.8)
5.54 5.65 5.48 5.64
(0.27) (0.42) (0.59) (0.28)
HPRS (TiB2–10 vol.% NiAl) L M N O (annealed)
900, 900, 900, 900,
15 30 45 30
98.3 98.2 98.2 97.4
20.4 21.6 21.6 20.9
(0.8) (1.5) (1.2) (0.6)
4.10 4.38 3.95 4.45
(0.21) (0.32) (0.57) (0.19)
RHP (TiB2–30 vol.% NiAl) P Q R
1650, 1 1650, 15 1650, 30
90 100 100
— 15.0 (1.2) 15.2 (1.1)
— 3.82 (0.27) 3.45 (0.31)
RHP (TiB2–15 vol.% NiAl) S T
1650, 15 1650, 30
100 100
19.9 (1.6) 20.4 (1.3)
3.05 (0.42) 3.12 (0.26)
RHP (TiB2–10 vol.% NiAl) U V
1650, 15 1650, 30
97.5 100
21.9 (1.0) 22.1 (1.2)
3.01 (0.35) 2.97 (0.28)
5.2 (0.3) 5.4 (0.3) 4.0 (0.3)
— — —
X is the mean value, s the standard deviation and %TD the percent theoretical density.
gas inside the pores expands resulting in the increase in porosity (Table 2; Fig. 1). NiAl compact prepared at 900°C contained NiAl, Ni2Al3, and Ni3Al phases (Fig. 2(a)). These phases were metallographically identified by etching the polished samples with Kallings solution-2. This solution reacts
faster with Ni3Al than either NiAl or Ni2Al3. The white phases are NiAl and Ni2Al3, whereas the grey phase is Ni3Al (Fig. 3(a)). The reaction sintering of NiAl consists of a complex reaction with several sequential steps. First the Al-rich compounds NiAl3 and Ni2Al3 are formed. Subsequently, these phases react with Ni to
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form NiAl. Depending on the availability of Ni in the vicinity of NiAl, Ni3Al can also form [25 – 27]. Therefore, in a reaction sintered compact where the reaction is incomplete, all the four intermetallic phases are likely to be present. In HPRS, the heat loss in the reaction zone by conduction is high leading to incomplete reaction. The reaction is completed when these compacts are annealed at 1100°C for 120 min (Fig. 2(b)). The microstructure of the annealed sample shows a single phase material (Fig. 3(b)). Nickel and aluminium are ductile and, therefore, deform significantly during pressurization. The material flow patterns are preserved even after the HPRS (Fig. 3(a)). The annealing treatment does not seem to change this grain morphology (Fig. 3(b)). The hardness values of the HPRS compacts (5.2–5.4 GPa) are higher than that of the NiAl ( 4 GPa) reported by Cheng et al. [16]. The high hardness is due to the high internal stress and presence of Ni2Al3 phase (hardness of 6.2 GPa). The hardness of the compact decreases to 4 GPa upon annealing, which is comparable to the hardness of NiAl.
Fig. 2. X-ray diffraction patterns of NiAl. (a) HPRS and (b) HPRS followed by annealing at 1100°C for 120 min. HPRS was conducted at 3 GPa and 900°C for 30 min. The compact contains various intermediate phases (Ni3Al and Ni2Al3). After annealing, the compact contains only NiAl phase.
3.2. TiB2 –NiAl composites (HPRS)
Fig. 1. Photomicrographs showing pore morphology and pore size. (a) HPRS and (b) HPRS followed by annealing at 1100°C for 120 min. The HPRS was conducted at 3 GPa and 900°C for 30 min. Annealing results in the increase in both porosity and pore size.
The densities of the TiB2 –NiAl composites prepared by HPRS are listed in Table 2. The porosity does not change significantly either on increasing the sintering time (from 15 to 45 min) or decreasing the amount of matrix NiAl phase (from 30 to 15 vol.%). These compacts contained residual porosity of 1%. A decrease in NiAl content to 10 vol.% results in the increase in porosity. The maximum density obtained in this case is 98%. Similar to NiAl, annealing of these compacts results in the increase in porsity and pore size. The densification mechanism in TiB2 –NiAl composites is different from that in pure NiAl. The major volume fraction in these composites is TiB2 phase (70– 90 vol.%). During pressurization, the densification takes place by the rearrangement of the TiB2 particles. The yield strength of Ni and Al being very low as compared to that of TiB2, these ductile phases deform significantly and fill the voids. On raising the temperature, the reaction starts and a liquid phase is formed. The volume fraction of the liquid phase decreases with the progress of the reaction. Finally, the liquid disappears when the reaction is complete. However, presence of the liquid even for a short period is beneficial for densification, as second stage of particle rearrangement takes place during this period. In the TiB2 –NiAl composites, the volume fraction of liquid phase and the heat generated due to reaction are less as compared to
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the pure NiAl. As a consequence, the time required for densification is longer. When the volume fraction of the matrix phase is low, the amount of liquid phase is not sufficient to fill all the pores resulting in the decrease in sintered density (Table 2). The XRD patterns (Fig. 4(a)) indicate that the composites contain mainly TiB2, NiAl and Ni3Al phases. After annealing, the Ni3Al disappears (Fig. 4(b)). The microstructures consist of TiB2 grains embedded in the matrix phase (Figs. 5 and 6). Due to the low synthesis temperature, the problem of grain growth is not encountered in HPRS. Irrespective of the composition and the synthesis time, all the compacts have similar microstructures. The mechanical properties of the composites are listed in Table 2. The hardness and fracture toughness values are in the range 10.4 – 21.6 GPa and 3.9–5.6 MPa m, respectively. TiB2 being the harder phase, the hardness of the composites increases with the decrease in the amount of the matrix NiAl phase. Increase in the sintering time from 15 to 30 min results
Fig. 3. Optical microstructures of NiAl. (a) HPRS and (b) HPRS followed by annealing at 1100°C for 120 min. The HPRS was conducted at 3 GPa and 900°C for 30 min. In (a) the grey phase is Ni3Al and white phases are NiAl and Ni2Al3. After annealing, the microstructure is uniform confirming a single phase material.
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in a better bonding between the matrix and the TiB2 phases leading to the improvement in mechanical properties. Further increase in time, causes grain growth in the matrix phase resulting in deterioration of the toughness. The improvement in fracture toughness after annealing is mainly due to the absence of brittle Ni2Al3 phase. Annealing also leads to homogenization of the matrix phase which results in less nonuniformity in the mechanical properties (Table 2). Fracture toughness values remain almost constant when the volume fraction of the matrix phase is reduced to 0.15 from 0.30. A further decrease to 10 vol.% results in compacts with inferior toughness.
3.3. TiB2 –NiAl composites (RHP) In RHP, though the reaction between the reactant phases were complete in 1 min (Fig. 7), the compacts contained a substantial amount of porosity (Table 2). Fully dense compacts were obtained only when the temperature was maintained for 15 min. In case of composites with 10 vol.% NiAl, a higher hot pressing time (30 min) was required to obtain full density. The RHP temperature (1650°C) is higher than the melting points of all the intermetallics in the Ni–Al system. This results in the formation of a liquid phase, and the densification proceeds through liquid phase sintering. The time required for complete densification depends on the volume fraction of the liquid phase. It increases with the decrease in the volume fraction of the liquid phase forming additives and vice-versa. The XRD patterns of the TiB2 –30 vol.%NiAl hot pressed for 1 and 30 min are shown in Fig. 7. In both the samples, no elemental Ni or Al was present. However, as compared to HPRS composites (Fig. 2), some additional phases (Ni2B, AlB2 and NiTi2) are detected in the RHP composites (Fig. 7). The formation of the liquid during RHP enhances the reaction betweenTiB2 and the other constituents of the composites. A part of Ni and Al is consumed to form Ni2B and AlB2 phases. The microstructures consist of equiaxed grains of TiB2 in the matrix of Ni–Al intermetallics (Fig. 8). The increase in hot pressing time results in grain growth. In RHP, the grain growth takes place by solution and reprecipitation. The smaller grains dissolve in the liquid and reprecipitate on the existing larger grains. Increase in temperature and time results in more material transport and enhances grain growth. Below a critical volume fraction of the liquid phase, where the contact area between TiB2 grains is large, grain growth by solid state diffusion is also appreciable. In general, the hot pressed composites have better hardness (15–22 GPa) and inferior toughness (2.97– 3.82 MPa m) than those of HPRS composites (Table 2). The differences in these properties are mainly due to the presence of Ni2B and AlB2 in RHP compacts. These
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Fig. 4. XRD patterns of TiB2 –30 vol.%NiAl composites after (a) HPRS and (b) HPRS followed by annealing at 1100°C for 120 min. The HPRS was conducted at 3 GPa and 900°C for 30 min. The composite contains mainly TiB2, NiAl and Ni3Al phases.
Fig. 5. SEM fractographs of HPRS TiB2 –NiAl composites. (a) 30 vol.%NiAl, (b) 15 vol.%NiAl and (c) 10 vol.%NiAl. The HPRS was conducted at 3 GPa and 900°C for 30 min. Microstructures consist of TiB2 grains embedded in the NiAl matrix.
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Fig. 6. SEM fractographs of HPRS TiB2 –30 vol.%NiAl composites. The HPRS was conducted at 3 GPa and 900°C for (a) 15 min and (b) 45 min. No grain growth is observed due to increase in synthesis time.
borides are formed at the cost of NiAl phase. These phases are harder and less ductile than Ni – Al intermetallic phases. Also, the microstructures of the RHP composites are much coarser than those fabricated by HPRS.
4. Conclusions TiB2 –NiAl (10 – 30 vol.%) composites with 99% density have been fabricated by reaction sintering at high pressure (3 GPa) and low temperature (900°C). Application of pressure is beneficial in obtaining high sintered density at low temperatures, but the associated heat transfer process is the key factor that controls the degree of reaction. The heat loss by conduction from the reaction zone leads to incomplete reaction. As a result, the compacts contain various intermediate phases (NiAl3, Ni2Al3). Annealing of the HPRS compacts is essential to enhance the reaction between these intermediate phases to obtain a single phase NiAl ma-
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Fig. 7. XRD patterns of TiB2 – 30 vol.%NiAl composites fabricated by RHP. Hot pressing was carried out at 20 MPa and 1650°C for (a) 1 min and (b) 30 min. The composites contain borides of Ni and Al in addition to the phases present in TiB2 – NiAl composites fabricated by HPRS.
trix. The composites fabricated have hardness in the range 10 to 20 GPa and fracture toughness in the range 3.5 to 5.6 MPa m. The RHP at 1650°C, results in the formation of Ni2B and AlB2 phases. These phases impart brittleness to the composites. The RHP composites have better hardness (15–22 GPa) but inferior toughness (2.9–3.8 MPa m) than those fabricated by HPRS.
Acknowledgements The authors thank. P.M. Jaman for the assistance rendered in the fabrication of the high pressure cells. The help of M.A. Venkatswamy and S. Usha Devi for electron microscopic and X-ray studies, respectively is acknowledged. This work was funded by BRNS, India (grant no.34/2/96R&D-II/1066).
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Fig. 8. SEM fractographs of TiB2 –15 vol.%NiAl composites fabricated by RHP. Hot pressing was carried out at 20 MPa and 1650°C for (a) 15 min and (b) 30 min. The microstructures reveal abnormal grain growth in TiB2 phase (gray). The microstructures are very coarse as compared to those processed by HPRS (Fig. 5).
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Reaction sintering of NiAl and TiB2 –NiAl composites under pressure S.K. Bhaumik, C. Divakar, L. Rangaraj, A.K. Singh * High Pressure Laboratory, Materials Science Di6ision, National Aerospace Laboratories, Bangalore 560 017, India Received 7 May 1998
Abstract Intermetallic matrix composites are a new class of engineering materials for high temperature structural applications in oxidizing and aggressive environments. Attempts have been made in the present investigation to synthesize TiB2 –NiAl composites, in which the matrix phase NiAl was produced in situ by reaction synthesis. The composites with 10, 15 and 30 vol.%NiAl were fabricated from the mixtures of elemental Ni and Al and TiB2 powders by high pressure reaction sintering (HPRS) and reactive hot pressing (RHP). The HPRS and RHP were carried out at 3 GPa and 900°C, and 20 MPa and 1650°C, respectively. The different phases in the sintered compacts were identified by X-ray diffraction and microstructural studies. In HPRS, the reaction was incomplete which gave rise to various intermediate phases (Ni2Al3, Ni3Al). It was necessary to anneal these compacts at 1100°C to obtain TiB2 –NiAl composites with single phase NiAl matrix. The densities of the HPRS compacts were 99%. The hardness and fracture toughness values were in the range 10 to 20 GPa and 3.9 to 5.7 MPa m, respectively. The RHP compacts contained AlB2, Ni2B and NiTi2 phases in addition to those present in the HPRS compacts. The RHP composites were fully dense. These had superior hardness (15 – 22 GPa) but inferior fracture toughness (2.9 – 3.8 MPa m) compared to those obtained by HPRS. © 1998 Elsevier Science S.A. All rights reserved. Keywords: Composites; Sintering; Matrix phase; High pressure
1. Introduction The unique combination of high melting point (2900°C), high hardness (27 GPa), high elastic modulus (370 GPa), high electrical conductivity, and relatively low coefficient of thermal expansion (4.6× 10 − 6°C) has made TiB2 an important material for high performance applications [1–4]. Attempts have been made [5–7] to consolidate TiB2 powder into dense solid bodies by reaction sintering, and hot pressing with or without sintering aids. In most cases, the consolidation by solid state sintering requires temperatures of 2100°C and times over 60 min. This often results in exaggerated grain growth leading to inferior mechanical properties. The sintering temperature and time can be reduced with the use of liquid phase forming additives. In such cases, * Corresponding author: Tel.: +91 80 5270098; fax: + 91 80 5260826; e-mail: [email protected]
a low melting phase is formed at the grain boundaries which adversely affects the high temperature mechanical properties of the products. Nickel aluminide containing more than 41 at.%Al forms a single phase ordered B2 structure based on the body centered cubic (bcc) lattice [8]. In terms of thermophysical properties, NiAl (B2) is more suitable for high temperature applications than Ni3Al (L12) [8,9]. It has high melting point (1638°C), low density (5.86 g cm − 3), high young’s modulus (294 GPa), and high thermal conductivity (76 Wm − 1 K − 1) at ambient temperature [10–12]. In addition, NiAl has excellent oxidation resistance at high temperatures (1000°C). However, the structural use of NiAl suffers from poor fracture resistance at ambient temperatures, and low strength and creep resistance at elevated temperatures. Various processing techniques such as rapid solidification [13], mechanical alloying [14], reaction sintering [11,15] and RHP are normally used for the synthesis of
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NiAl intermetallic. HPRS of stoichiometric and Ni-rich NiAl has also been reported recently by Cheng et al. [16]. The Ni–49.3 at.%Al produced by HPRS at 2 GPa and 500°C is reported [16] to have high density (98.2%)), good compressive ductility (14.5%), and strength (876 MPa). Because of the good high temperature properties of NiAl, efforts have been made to use this intermetallic as a matrix for high temperature composites. Various composites of NiAl with Al2O3, Y2O3 and TiB2 have been reported [11,17 – 20]. The effects of fine TiB2 (1–3 mm) particles on NiAl have been studied in detail [17–20]. In these studies, the addition of the refractory particulate was limited to maximum 30 vol.%. These particles improve the compressive strength of stoichiometric NiAl significantly. In the present investigation,TiB2 – NiAl (10–30 vol.%) composites have been prepared by HPRS and RHP. The stoichiometric composition of NiAl has been chosen as the matrix phase. The results are presented in this paper.
The RHP was conducted at 1650°C in flowing argon atmosphere for 1–30 min. The heating was done by an induction furnace. The heating rate was 15°C s − 1. After HPRS and RHP, the reaction products were identified by X-ray diffraction. The density was measured by volume displacement method. The theoretical density used for comparison was calculated by the volumetric rule of mixtures assuming that Ni and Al reacted completely to form NiAl. The microstructural observations were carried out in optical (Neophot-2) and JEOL 840A scanning electron microscopes (SEM). Shimadzu HSV-20 microhardness tester with a load of 1.96 N was used to determine the hardness. For indentation fracture toughness determination, a load of 49 N was used. The fracture toughness was calculated from the equation [23]: KIC = 0.0824 (P/C 3/2) where, P is the indentation load (N) and C the crack length (mm). Ten indentations were made on each sample for hardness and fracture toughness measurements, and the average values are reported.
2. Experimental 3. Results and discussion The raw materials used in the present study were elemental nickel and aluminium, and TiB2 powders. A master alloy powder mixture of Ni – 50 at.%Al was first prepared by ball milling (Fritsch Pulverisette-5 centrifugal mill) in hexane medium for 24 h. This master alloy powder was used to prepare the composite powder mixtures of desired compositions. In the second step, the milling was carried out for 8 h. The details of the powder characteristics are given in Table 1. The powder mixture compositions and the synthesis conditions are given in Table 2. The HPRS of NiAl and TiB2 – NiAl composites was carried out in a 200 t cubic anvil apparatus (anvil face, 10 mm square) capable of generating pressures up to 6.5 GPa and temperatures up to 3000°C. The details of the high pressure-high temperature cells used in these experiments are reported elsewhere [21]. In an actual experiment, the sample was first pressurized to 3 GPa, and then heated to desired temperature (800 –900°C). The heating rate was 40°C s − 1. The temperatures were maintained for 15 – 45 min. The typical dimensions of the recovered samples were 3.5 mm in diameter and 4 mm in thickness. Some of these compacts were annealed at 1100°C in vacuum (10 − 4 Torr) for 120 min. The RHP of TiB2 – NiAl composites was carried out in a hot press fabricated in the laboratory [22]. The powder mixture was cold compacted at 400 MPa into a cylindrical compact of 13 mm diameter and 8 mm height. These compacts, with green densities in the range 55 to 60% were placed in a graphite die and a load corresponding to pressure of 20 MPa was applied.
3.1. Matrix phase NiAl (HPRS) The measured densities (as a percentage of theoretical density) of HPRS NiAl compacts fabricated at 800 and 900°C for 15 min are listed in Table 2. The maximum density obtained is 99.3%. During HPRS, the exothermic reaction between Ni and Al generates enough heat to produce a liquid phase. This liquid provides a faster diffusion than encountered in the solid state leading to enhanced densification. The amount of liquid phase formed depends on the powder composition (Al content) and temperature. It has been estimated [24] that stoichiometric Ni–Al composition can produce as much as 50 vol.% liquid when the temperature is in the range of 600–700°C. The application of pressure during reaction sintering helps in eliminating the big pores often observed in the pressureless sintering. It has been reported [25] that during the reaction synthesis of various intermetallics in the Ni–Al system, a pressure of 100 MPa is sufficient to obtain 99.9% dense compacts. In the present investigation, even though the pressure was much higher (3 GPa), the sintered compacts had 1% porosity. This porosity is due to the adsorbed gases and volatile impurities in the reactant powder mixtures. During synthesis, the released gaseous phases fill the pores and hinder complete densification. The high pressure cell is completely isolated from the atmosphere and the trapped gases can not escape from the compact. When the HPRS compacts are annealed at higher temperature (1100°C), the
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Table 1 Powder characteristicsa Material
TiB2
Nickel
Aluminium
Source
HC Starck, Germany
INCO (UK)
ALCOA
Chemical analysis
B (30.3) C (0.3) O (0.25) N (0.39) Ti (Balance)
C (0.06) Fe (0.005) O (0.05) Co (0.0003) N2 (0.03) S (0.003) Ni (Balance)
High purity atomized powder 7123
10
3.7
16.5
Particle size (mm) a
Al (99.3)
All compositions given in wt.%.
Table 2 Density, hardness and fracture toughness values of NiAl and TiB2–NiAl (10–30 vol.%) composites fabricated by HPRS and RHP Sample
Synthesis temperature (°C) time (min)
Density (% TD)
Hardness (Hv) (GPa) X (s)
X (s)
HPRS (NiAl) A B C (annealed)
800, 15 900, 15 900, 15
99.1 99.3 97.5
HPRS (TiB2–30 vol.%NiAl) D E F G (annealed)
900, 900, 900, 900,
15 30 40 30
99.2 99.1 99.1 98.7
10.5 13.2 14.2 12.8
(1.6) (0.6) (1.5) (0.6)
5.10 5.69 5.43 5.53
(0.60) (0.28) (0.41) (0.28)
HPRS (TiB2–15 vol.% NiAl) H I J K (annealed)
900, 900, 900, 900,
15 30 45 30
99.2 99.0 99.3 98.8
16.4 17.8 19.3 18.2
(1.4) (2.0) (1.1) (0.8)
5.54 5.65 5.48 5.64
(0.27) (0.42) (0.59) (0.28)
HPRS (TiB2–10 vol.% NiAl) L M N O (annealed)
900, 900, 900, 900,
15 30 45 30
98.3 98.2 98.2 97.4
20.4 21.6 21.6 20.9
(0.8) (1.5) (1.2) (0.6)
4.10 4.38 3.95 4.45
(0.21) (0.32) (0.57) (0.19)
RHP (TiB2–30 vol.% NiAl) P Q R
1650, 1 1650, 15 1650, 30
90 100 100
— 15.0 (1.2) 15.2 (1.1)
— 3.82 (0.27) 3.45 (0.31)
RHP (TiB2–15 vol.% NiAl) S T
1650, 15 1650, 30
100 100
19.9 (1.6) 20.4 (1.3)
3.05 (0.42) 3.12 (0.26)
RHP (TiB2–10 vol.% NiAl) U V
1650, 15 1650, 30
97.5 100
21.9 (1.0) 22.1 (1.2)
3.01 (0.35) 2.97 (0.28)
5.2 (0.3) 5.4 (0.3) 4.0 (0.3)
— — —
X is the mean value, s the standard deviation and %TD the percent theoretical density.
gas inside the pores expands resulting in the increase in porosity (Table 2; Fig. 1). NiAl compact prepared at 900°C contained NiAl, Ni2Al3, and Ni3Al phases (Fig. 2(a)). These phases were metallographically identified by etching the polished samples with Kallings solution-2. This solution reacts
faster with Ni3Al than either NiAl or Ni2Al3. The white phases are NiAl and Ni2Al3, whereas the grey phase is Ni3Al (Fig. 3(a)). The reaction sintering of NiAl consists of a complex reaction with several sequential steps. First the Al-rich compounds NiAl3 and Ni2Al3 are formed. Subsequently, these phases react with Ni to
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form NiAl. Depending on the availability of Ni in the vicinity of NiAl, Ni3Al can also form [25 – 27]. Therefore, in a reaction sintered compact where the reaction is incomplete, all the four intermetallic phases are likely to be present. In HPRS, the heat loss in the reaction zone by conduction is high leading to incomplete reaction. The reaction is completed when these compacts are annealed at 1100°C for 120 min (Fig. 2(b)). The microstructure of the annealed sample shows a single phase material (Fig. 3(b)). Nickel and aluminium are ductile and, therefore, deform significantly during pressurization. The material flow patterns are preserved even after the HPRS (Fig. 3(a)). The annealing treatment does not seem to change this grain morphology (Fig. 3(b)). The hardness values of the HPRS compacts (5.2–5.4 GPa) are higher than that of the NiAl ( 4 GPa) reported by Cheng et al. [16]. The high hardness is due to the high internal stress and presence of Ni2Al3 phase (hardness of 6.2 GPa). The hardness of the compact decreases to 4 GPa upon annealing, which is comparable to the hardness of NiAl.
Fig. 2. X-ray diffraction patterns of NiAl. (a) HPRS and (b) HPRS followed by annealing at 1100°C for 120 min. HPRS was conducted at 3 GPa and 900°C for 30 min. The compact contains various intermediate phases (Ni3Al and Ni2Al3). After annealing, the compact contains only NiAl phase.
3.2. TiB2 –NiAl composites (HPRS)
Fig. 1. Photomicrographs showing pore morphology and pore size. (a) HPRS and (b) HPRS followed by annealing at 1100°C for 120 min. The HPRS was conducted at 3 GPa and 900°C for 30 min. Annealing results in the increase in both porosity and pore size.
The densities of the TiB2 –NiAl composites prepared by HPRS are listed in Table 2. The porosity does not change significantly either on increasing the sintering time (from 15 to 45 min) or decreasing the amount of matrix NiAl phase (from 30 to 15 vol.%). These compacts contained residual porosity of 1%. A decrease in NiAl content to 10 vol.% results in the increase in porosity. The maximum density obtained in this case is 98%. Similar to NiAl, annealing of these compacts results in the increase in porsity and pore size. The densification mechanism in TiB2 –NiAl composites is different from that in pure NiAl. The major volume fraction in these composites is TiB2 phase (70– 90 vol.%). During pressurization, the densification takes place by the rearrangement of the TiB2 particles. The yield strength of Ni and Al being very low as compared to that of TiB2, these ductile phases deform significantly and fill the voids. On raising the temperature, the reaction starts and a liquid phase is formed. The volume fraction of the liquid phase decreases with the progress of the reaction. Finally, the liquid disappears when the reaction is complete. However, presence of the liquid even for a short period is beneficial for densification, as second stage of particle rearrangement takes place during this period. In the TiB2 –NiAl composites, the volume fraction of liquid phase and the heat generated due to reaction are less as compared to
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the pure NiAl. As a consequence, the time required for densification is longer. When the volume fraction of the matrix phase is low, the amount of liquid phase is not sufficient to fill all the pores resulting in the decrease in sintered density (Table 2). The XRD patterns (Fig. 4(a)) indicate that the composites contain mainly TiB2, NiAl and Ni3Al phases. After annealing, the Ni3Al disappears (Fig. 4(b)). The microstructures consist of TiB2 grains embedded in the matrix phase (Figs. 5 and 6). Due to the low synthesis temperature, the problem of grain growth is not encountered in HPRS. Irrespective of the composition and the synthesis time, all the compacts have similar microstructures. The mechanical properties of the composites are listed in Table 2. The hardness and fracture toughness values are in the range 10.4 – 21.6 GPa and 3.9–5.6 MPa m, respectively. TiB2 being the harder phase, the hardness of the composites increases with the decrease in the amount of the matrix NiAl phase. Increase in the sintering time from 15 to 30 min results
Fig. 3. Optical microstructures of NiAl. (a) HPRS and (b) HPRS followed by annealing at 1100°C for 120 min. The HPRS was conducted at 3 GPa and 900°C for 30 min. In (a) the grey phase is Ni3Al and white phases are NiAl and Ni2Al3. After annealing, the microstructure is uniform confirming a single phase material.
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in a better bonding between the matrix and the TiB2 phases leading to the improvement in mechanical properties. Further increase in time, causes grain growth in the matrix phase resulting in deterioration of the toughness. The improvement in fracture toughness after annealing is mainly due to the absence of brittle Ni2Al3 phase. Annealing also leads to homogenization of the matrix phase which results in less nonuniformity in the mechanical properties (Table 2). Fracture toughness values remain almost constant when the volume fraction of the matrix phase is reduced to 0.15 from 0.30. A further decrease to 10 vol.% results in compacts with inferior toughness.
3.3. TiB2 –NiAl composites (RHP) In RHP, though the reaction between the reactant phases were complete in 1 min (Fig. 7), the compacts contained a substantial amount of porosity (Table 2). Fully dense compacts were obtained only when the temperature was maintained for 15 min. In case of composites with 10 vol.% NiAl, a higher hot pressing time (30 min) was required to obtain full density. The RHP temperature (1650°C) is higher than the melting points of all the intermetallics in the Ni–Al system. This results in the formation of a liquid phase, and the densification proceeds through liquid phase sintering. The time required for complete densification depends on the volume fraction of the liquid phase. It increases with the decrease in the volume fraction of the liquid phase forming additives and vice-versa. The XRD patterns of the TiB2 –30 vol.%NiAl hot pressed for 1 and 30 min are shown in Fig. 7. In both the samples, no elemental Ni or Al was present. However, as compared to HPRS composites (Fig. 2), some additional phases (Ni2B, AlB2 and NiTi2) are detected in the RHP composites (Fig. 7). The formation of the liquid during RHP enhances the reaction betweenTiB2 and the other constituents of the composites. A part of Ni and Al is consumed to form Ni2B and AlB2 phases. The microstructures consist of equiaxed grains of TiB2 in the matrix of Ni–Al intermetallics (Fig. 8). The increase in hot pressing time results in grain growth. In RHP, the grain growth takes place by solution and reprecipitation. The smaller grains dissolve in the liquid and reprecipitate on the existing larger grains. Increase in temperature and time results in more material transport and enhances grain growth. Below a critical volume fraction of the liquid phase, where the contact area between TiB2 grains is large, grain growth by solid state diffusion is also appreciable. In general, the hot pressed composites have better hardness (15–22 GPa) and inferior toughness (2.97– 3.82 MPa m) than those of HPRS composites (Table 2). The differences in these properties are mainly due to the presence of Ni2B and AlB2 in RHP compacts. These
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Fig. 4. XRD patterns of TiB2 –30 vol.%NiAl composites after (a) HPRS and (b) HPRS followed by annealing at 1100°C for 120 min. The HPRS was conducted at 3 GPa and 900°C for 30 min. The composite contains mainly TiB2, NiAl and Ni3Al phases.
Fig. 5. SEM fractographs of HPRS TiB2 –NiAl composites. (a) 30 vol.%NiAl, (b) 15 vol.%NiAl and (c) 10 vol.%NiAl. The HPRS was conducted at 3 GPa and 900°C for 30 min. Microstructures consist of TiB2 grains embedded in the NiAl matrix.
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Fig. 6. SEM fractographs of HPRS TiB2 –30 vol.%NiAl composites. The HPRS was conducted at 3 GPa and 900°C for (a) 15 min and (b) 45 min. No grain growth is observed due to increase in synthesis time.
borides are formed at the cost of NiAl phase. These phases are harder and less ductile than Ni – Al intermetallic phases. Also, the microstructures of the RHP composites are much coarser than those fabricated by HPRS.
4. Conclusions TiB2 –NiAl (10 – 30 vol.%) composites with 99% density have been fabricated by reaction sintering at high pressure (3 GPa) and low temperature (900°C). Application of pressure is beneficial in obtaining high sintered density at low temperatures, but the associated heat transfer process is the key factor that controls the degree of reaction. The heat loss by conduction from the reaction zone leads to incomplete reaction. As a result, the compacts contain various intermediate phases (NiAl3, Ni2Al3). Annealing of the HPRS compacts is essential to enhance the reaction between these intermediate phases to obtain a single phase NiAl ma-
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Fig. 7. XRD patterns of TiB2 – 30 vol.%NiAl composites fabricated by RHP. Hot pressing was carried out at 20 MPa and 1650°C for (a) 1 min and (b) 30 min. The composites contain borides of Ni and Al in addition to the phases present in TiB2 – NiAl composites fabricated by HPRS.
trix. The composites fabricated have hardness in the range 10 to 20 GPa and fracture toughness in the range 3.5 to 5.6 MPa m. The RHP at 1650°C, results in the formation of Ni2B and AlB2 phases. These phases impart brittleness to the composites. The RHP composites have better hardness (15–22 GPa) but inferior toughness (2.9–3.8 MPa m) than those fabricated by HPRS.
Acknowledgements The authors thank. P.M. Jaman for the assistance rendered in the fabrication of the high pressure cells. The help of M.A. Venkatswamy and S. Usha Devi for electron microscopic and X-ray studies, respectively is acknowledged. This work was funded by BRNS, India (grant no.34/2/96R&D-II/1066).
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Fig. 8. SEM fractographs of TiB2 –15 vol.%NiAl composites fabricated by RHP. Hot pressing was carried out at 20 MPa and 1650°C for (a) 15 min and (b) 30 min. The microstructures reveal abnormal grain growth in TiB2 phase (gray). The microstructures are very coarse as compared to those processed by HPRS (Fig. 5).
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