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Effect of NbC addition on mechanical properties of dual two-phase Ni3Al–Ni3V intermetallic alloy
Materials Science and Engineering A 527 (2010) 6012–6019
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Effect of NbC addition on mechanical properties of dual two-phase Ni3 Al–Ni3 V intermetallic alloy Y. Kitaura, Y. Kaneno, T. Takasugi ∗ Department of Materials Science, Osaka Prefecture University, 1-1 Gakuen-cho, Naka-ku, Sakai, Osaka, 599-8531, Japan
a r t i c l e
i n f o
Article history: Received 4 February 2010 Received in revised form 28 May 2010 Accepted 4 June 2010
Keywords: Multiphase intermetallics Mechanical properties at high temperatures Heat treatment Microstructure NbC carbide
a b s t r a c t Dual two-phase intermetallic alloy that has an alloy composition of Ni75 Al9 V13 Nb3 and is composed of geometrically close packed (GCP) Ni3 Al (L12 ) and Ni3 V (D022 ) phases was studied, focusing on the effect of NbC addition on high-temperature tensile properties. The two-phase microstructures defined by primary Ni3 Al precipitates and eutectoid region consisting of Ni3 Al and Ni3 V phases were kept and were dispersed by Nb2 C particles regardless of NbC content. EPMA analysis showed that alloy compositions in matrix of the NbC added alloy were different from those in the base alloy. The positive temperature dependence of flow strength was observed for all NbC added alloys. In low temperature range the maximum strengthening took place at 1 at.% NbC content. On the other hand, tensile elongation increased with increasing NbC content up to 1 at.% NbC content and gradually decreased at NbC content over 1 at.% content in whole temperature ranges. Also, the NbC addition resulted in change of the fracture mode from brittle transgranular fracture to ductile transgranular fracture in low temperature range, and from brittle intergranular fracture to ductile transgranular in high temperature range. Possible mechanisms responsible for the strengthening and ductilization by NbC addition were discussed, based on the behavior of C solutes in the matrix, which were released from the charged NbC carbides. © 2010 Elsevier B.V. All rights reserved.
1. Introduction Geometrically close packed (GCP) Ni3 X intermetallic phases generally exhibit high phase and microstructural stabilities up to their melting points because of low atomic diffusivity owing to their close packed structures [1–4]. Multi-phase intermetallic alloys based on GCP Ni3 X structures have been studied for various Ni3 X (X: Al, Ti and Nb [5,6], X: Si, Ti and Nb [7–9], X: Al, Ti and V [10–14], and X: Al, Nb and V [15–18]) alloy systems to develop a new type of high-temperature structural materials. Among so far studied alloy systems, so-called dual two-phase intermetallic alloys based on Ni3 Al (L12 )–Ni3 V (D022 ) pseudo-binary alloy system were shown to be attractive specifically as high temperature structural materials. The attractive mechanical properties arise from a peculiar microstructure consisting of upper two-phase microstructure with a micron scale and lower two-phase microstructure with a sub-micron scale. The upper two-phase microstructure is composed of primary Ni3 Al precipitates and Ni solid solution (A1) (channel region) at high temperature, and the prior A1 phase is decomposed into Ni3 Al + Ni3 V by a eutectoid reaction at low temperature, resulting in the lower two-phase microstructure. The dual two-phase microstructures exhibit coherent and fine
∗ Corresponding author. Tel.: +81 72 254 9314; fax: +81 72 254 9912. E-mail address: [email protected] (T. Takasugi). 0921-5093/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2010.06.014
microstructure and high microstructural stability [10,11,13,14]. In addition, high temperature mechanical properties (such as tensile and creep strength) of the dual two-phase intermetallic alloys were found to be superior to those of many conventional super alloys [12,14,18]. Considering the development of the dual two-phase intermetallic alloys as next generation-type high temperature structural materials, it is necessary to furthermore improve their mechanical properties. Along this strategy, traditional strengthening or toughening mechanisms that have been employed to many conventional steels and Ni alloys should be applied. Actually, Ti [12,14] and Nb [16–18] resulting in solid solution strengthening have been so far employed to the dual two-phase intermetallic alloys, and also boron [12,16,17] improving tensile elongation has been doped into the dual two-phase intermetallic alloys. In this study, expecting dispersion strengthening due to Nb carbides and/or solid solution strengthening due to C atoms released from Nb carbides, NbC carbides ranging to 5.0 at.% were charged to the dual two-phase intermetallic alloys composed of Ni3 Al and Ni3 V phases containing Nb and doped with boron. Microstructure, room-temperature hardness, and tensile properties of the dual two-phase intermetallic alloys were investigated as functions of NbC content and test temperature. Improved tensile strength and elongation were observed in certain levels of NbC content and in a wide range of test temperature. Strengthening and ductilization mechanisms due to Nb carbides were discussed, based on the experimental results.
Y. Kitaura et al. / Materials Science and Engineering A 527 (2010) 6012–6019 Table 1 Alloy compositions used in this study (at.%).
Base 0.2 NbC 0.5 NbC 1.0 NbC 2.5 NbC 5.0 NbC a
Ni
Al
V
Nb
NbCa
Ba (wt ppm)
75 75 75 75 75 75
9 9 9 9 9 9
13 13 13 13 13 13
3 3 3 3 3 3
– 0.2 0.5 1.0 2.5 5.0
100 100 100 100 100 100
NbC and B are an extra number ; the numbers are not included in the total.
2. Experimental procedures Alloy buttons with a diameter of 50 mm were prepared by arc melting in argon gas atmosphere using tungsten electrode and copper hearth. The raw materials used in this study were 99.9 wt.% Ni, 99.99 wt.% Al, 99.9 wt.% V, 99.9 wt.% Nb, 99.9 wt.% boron (B) and 99 wt.% NbC. Boron was doped to suppress grain boundary fracture of the present intermetallic alloy [12,16,17]. For NbC carbides, powders with particle sizes of 1–3 m were used. Each alloy button was remelted more than three times to ensure chemical homogeneity through an entire button cross section. A composition expressed by Ni75 Al9 V13 Nb3 (at.%) doped with 100 wt ppm B is a base alloy composition adopted in this study. NbC was charged up to 5.0 at.% to the base alloy composition, as shown in Table 1. Hereafter, these alloy compositions are referred as base, 0.2 NbC, 0.5 NbC, 1.0 NbC, 2.5 NbC, and 5.0 NbC alloys in the followings. These alloys were homogenized at 1553 K for 3 h in a dynamic vacuum condition and then rapidly furnace cooled to room temperature to obtain the dual two-phase microstructure. Vickers hardness test was conducted in conditions of a holding time of 20 s and a load of 1 kg, using the plate-like specimens with a thickness of 1 mm that were sliced from the arc melted buttons. Hardness data points more than ten were collected and averaged in each experimental condition. Microstructures and alloy compositions in matrix and dispersions of the NbC added alloys were investigated by optical microscopy (OM) and scanning electron microscopy (JEOL: JXA870) with a wavelength-dispersive spectroscopy (WDS). The specimens were sliced from the heat-treated arc buttons along a solidified direction using an electro discharge machine (EDM). The sliced specimens were mechanically abraded on SiC paper and then electronically polished in a mixed solution of 15 ml H2 SO4 + 85 ml CH3 OH at 243 K. Also, X-ray diffraction (Rigaku: RINT-2500HKLC) was performed to determine constituent phases and second-phases, using Ni-filtrated Cu K␣ radiation at an accelerated voltage of 30 kV.
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Tensile specimens with a gauge dimension of 2 mm × 1 mm × 10 mm were cut from the heat-treated arc buttons using an EDM. Tensile axis of the tensile specimen was chosen in perpendicular to solidification direction of the arc button. The prepared tensile specimens were composed of bamboo-type grains in which most grain boundary planes are aligned in perpendicular to the tensile axis. Tensile tests were conducted in a temperature range between room temperature and 1173 K within a metal tube surrounded by an electronic furnace. The electro furnace was evacuated to a vacuum degree of approximately 1.5 × 10−3 Pa before testing. The tensile specimens were deformed at a strain rate of 1.66 × 10−4 s−1 up to fracturing. Tensile elongation was calculated from load-elongation curve drawn in the recorded chart, and defined as plastic strain to fracture. The fracture surface of the tensile deformed specimens was examined by using SEM. 3. Results 3.1. Microstructure Large columnar grains that were preferentially elongated along solidification direction and have a grain size of several mm were observed in arc melted material of the base alloy. The addition of NbC exceeding 0.2 at.% resulted in grain refinement and the grain size decreased from several mm to several hundred m with increasing NbC content as clearly shown in Fig. 1. The columnar grains similar to the base alloy were observed in all NbC added alloys. Fig. 2 shows SEM images of microstructures of the alloys observed at a low magnification. The base alloy and 0.2 NbC alloy exhibited a featureless microstructure consisiting of Ni3 Al + Ni3 V although two phases are not resolved in this SEM image. In the NbC added alloys (0.2 NbC, 0.5 NbC, 1.0 NbC and 5.0 NbC alloys are representatively shown in Fig. 2), there were a number of large bright Nb carbides with size of several m within grain interiors. It is found from these pictures that the size and density of the Nb carbides tend to increase with increasing NbC content, and also the Nb carbides tend to coalesce with increasing NbC content. It is suggested that the Nb carbides are formed as a primary phase and present in dendritic core region during solidification. X-ray diffraction patterns of base, 0.2 NbC, 0.5 NbC, 1.0 NbC, and 5.0 NbC alloys are shown in Fig. 3. In the base alloy, only diffraction lines indexed by both Ni3 Al and Ni3 V were identified. In the other NbC added alloys, additional diffraction lines indexed by NbC were identified and their lines were counted at higher angles than those calculated from a stoichiometric NbC [19,20]. This result reveals that the Nb carbides consisting of low C content, in other words, the Nb carbides to be assumed as the Nb2 C carbides are contained in
Fig. 1. Cross-section views of arc melted buttons of the (a) base, (b) 0.2 NbC, (c) 1.0 NbC, and (d) 5.0 NbC alloys.
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Fig. 2. SEM images of microstructures of the (a) base, (b) 0.2 NbC, (c) 0.5 NbC, (d) 1.0 NbC, and (e) 5.0 NbC alloys. Pictures were taken at a low magnification to show grain size and morphology, and NbC dispersions.
Fig. 3. X-ray diffraction patterns of the (a) base, (b) 0.2 NbC, (c) 0.5 NbC, (d) 1.0 NbC, and (e) 5.0 NbC alloys.
the dual two-phase intermetallic alloys, consistent with SEM-EPMA result shown later. SEM-EPMA results in which alloy compositions in matrix and Nb carbides of the base and 5.0 NbC alloys are shown in Table 2. Here, it is noted that the spot size for the EPMA observation for the matrix was large enough to examine average composition of the region decomposed into Ni3 Al and Ni3 V while that for the carbides was reduced as small as possible. Also for the analysis of carbon content, the standard specimen was used for obtaining accurate results. Alloy composition in the base alloy was almost equal to the nominal alloy composition, i.e., the intended alloy composition Table 2 Alloy compositions (at.%) in matrix and carbide in the base and 5.0 NbC alloys. Base
Ni
Al
V
Nb
Matrix
74.8
8.6
12.3
3.4
C –
5.0 NbC
Ni
Al
V
Nb
C
Matrix Carbide
74.1 2.43
9.4 0.016
11.3 8.66
3.8 60.9
1.42 28.0
shown in Table 1. In the Nb carbides in the 5.0 NbC added alloy, C content is much reduced from NbC. According to the composition of the Nb carbide it seems to be not NbC but Nb2 C. However, detailed analysis was not conducted in the crystal structure of the Nb carbides. Also, it was found that V atoms which are constituent element in the base alloy are contained in the Nb carbides. On the other hand, C atoms and surplus Nb atoms by about 0.8 at.%, which are assumed to be released from NbC carbides, are contained in the matrix in the 5.0 NbC added alloy. SEM images of microstructures in grain interior (matrix) of the alloys observed at a high magnification are shown in Fig. 4. The upper two-phase microstructures consisting of primary Ni3 Al precipitates and channel region is clearly imaged in each alloy. The size and morphology of the upper two-phase microstructures are similar and insensitive to NbC content. For the lower two-phase microstructure, i.e., the eutectoid region (channel region) corresponding to the base alloy composition, an extensive study has been done [16,17]. The eutectoid region, i.e. the prior Ni solid solution (A1) was shown to be composed of lamellar-like structures consisting of Ni3 Al and Ni3 V.
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Fig. 4. SEM images of microstructures of the (a) base, (b) 0.2 NbC, (c) 1.0 NbC, and (d) 5.0 NbC alloys. Pictures were taken at a high magnification to show the dual two-phase microstructure of each alloy.
3.2. Mechanical properties Change of room-temperature Vickers hardness with NbC content is shown in Fig. 5. As NbC content increases, the hardness monotonously increases up to about 2 at.% NbC content and then appears to be saturated in NbC content range beyond 2 at.%. Yield strength, tensile strength (UTS) and tensile elongation for the base, 0.2 NbC, 0.5 NbC, 1.0 NbC, 2.5 NbC, and 5.0 NbC alloys are shown in Fig. 6, respectively, as a function of temperature. In all alloys, positive temperature dependence of the yield strength took place, and at a peak temperature (around 900 K), the tensile strength as well as the yield strength almost exceeded 1 GPa for all alloys. The flow strength increase by NbC addition generally occurred in low temperature range below a peak temperature but little occurred in high temperature range beyond a peak temperature. The details will be described later. On the other hand, the improvement of tensile elongation by NbC addition was observed for all test temperatures. The details also will be described later.
Fig. 5. Change of Vickers hardness by NbC content.
Variations of yield strength, tensile strength (UTS), and tensile elongation with NbC content at room temperature (RT), 873 K, 1073 K, and 1173 K are shown in Fig. 7, respectively. It was clearly observed that in low temperature range (e.g. RT and 873 K) below a peak temperature, the yield strength as well as the tensile strength rapidly increased with increasing NbC content and made a peak at around 1.0 at.% NbC content, followed by a steady decrease (in the case of the tensile strength) or by consistently high level (in the case of the yield strength) in high NbC content range. The yield strength increase by about 450 MPa was thus observed at room temperature by the addition of 1 at.% NbC (Fig. 7). In high temperature range (e.g. 1073 K and 1173 K) beyond a peak temperature, the flow strength was little affected by NbC addition in their whole contents. Looking the variation curves of tensile elongation vs. NbC content, the tensile elongation increased with increasing NbC content and made a peak at around 1.0 at.% NbC content, followed by a decrease in high NbC content range. In the base alloy, preferential fracture mode at low temperature (below a peak temperature) was transgranular fracture accompanied by cleavage-like pattern while that at high temperature (beyond a peak temperature) was intergranular fracture mode accompanied by grain boundary facets on which tinny dimple patterns were observed, as shown in Figs. 8 and 9. However, fracture mode was apparently changed by NbC addition. That is, ductile transgranular fracture mode accompanied by higher fraction of dimple patterns was prevailing the fracture pattern in the entire deformation temperature tested (not only low temperature but also high temperature), as shown in Figs. 8 and 9. The observed fracture patterns reveal that largely energy-dissipated event occurs during propagation of micro-crack in the NbC added alloys not only at low temperature but also at high temperature, consequently improving the tensile elongation. Thus, the tensile behavior is consistent with the observed fractography: once grain boundary fracturing is suppressed, high tensile elongation which the dual two-phase microstructure intrinsically has can be obtained, resulting in the improvement of tensile ductility. However, large Nb carbides may act as the initiation sites of fracturing and may reduce tensile ductility of the present alloys. Fig. 10 shows the fractography of the
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Fig. 6. Yield strength, tensile strength (UTS), and tensile elongation as a function of test temperature for the (a) base, (b) 0.2 NbC, (c) 0.5 NbC, (d) 1.0 NbC, (e) 2.5 NbC, and (f) 5.0 NbC alloys, respectively.
5.0 at.% NbC added alloy in which preferential fracturing takes place from the Nb carbides themselves or the interface between Nb carbide and the matrix.
clarify precipitation phenomenon of the Nb carbides in solid state is an important issue to more sophistically control the dispersion of the Nb carbide particles and to improve mechanical properties of the dual two-phase intermetallic alloys.
4. Discussion 4.1. Strengthening by NbC addition It is assumed that the charged NbC carbides were wholly melted into liquid phase during arc melting and redeposited during solidification process although NbC with a cubic structure has very high formation enthalpy energy and therefore a very high melting point (3898 K) [19,20]. The result that the Nb carbides analyzed by EPMA showed different chemical composition from the charged NbC carbides, and at the same time surplus Nb atoms and C atoms were detected from the matrix of the NbC added alloys (Table 2) supports this speculation. Also, the morphology and size of the Nb carbides observed in the microstructures were apparently different from those of the charged NbC carbides. It is also speculated from SEM-EPMA (Fig. 2 and Table 2) and XRD (Fig. 3) results that the addition of NbC contents exceeding 0.5–1 at.% is high enough resulting in C solute atoms beyond the solubility limit (∼2 at.%) in the matrix of the present alloy. This value is not so much different from the reported solubility limit (several at.%) of C atoms in the major constituent phase Ni3 Al in the present alloy (composition) [21]. However, the lattice expansion due to C atoms soluble in Ni3 Al and/or Ni3 V was not correctly evaluated from X-ray diffraction patterns because the majorities of diffraction lines from both phases were superimposed and therefore are ambiguous and too broad. To
The strength properties observed in the base alloy basically are interpreted by the interfacial strengthening in addition to the strengthening due to two constituent phases [18]. The latter strengthening was basically interpreted in terms of the mixture rule of strength of two constituent phases [18]. The yield strength anomaly shown in Fig. 6 primarily arises from the nature that Ni3 Al and Ni3 V display the positive temperature dependence of the flow strength [4,22–24]. In both phases, the anisotropy in anti-phase boundary (APB) energy between the glide plane and non-glide plane promotes thermally activated cross-slip of the dissociated screw dislocation from the glide plane to non-glide plane, resulting in the strength anomaly. Besides the strength anomaly, large interfacial strengthening occurs in the present dual two-phase microstructures when a dislocation passes through the interface between the Ni3 Al and Ni3 V phases or between different variants in the Ni3 V phase [18]. Regarding mechanisms responsible for the additional strengthening due to NbC addition, the strengthening mechanism based on grain refinement is at least unlikely to be applied because the observed flow strength did not monotonously increase as
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Fig. 7. Variation of yield strength, tensile strength (UTS), and tensile elongation with NbC content at (a) room temperature (RT), (b) 873 K, (c) 1073 K, and (d) 1173 K, respectively.
grain size decreases (i.e., as NbC content increases). Also, the dispersion strengthening due to the Nb carbides is not applicable. A qualitative evaluation assuming the Orowan-type mechanism [25] predicts flow strength increase of several 10 MPa, apparently much lower than the strength increase (∼450 MPa) observed in the present study. Instead, the strengthening by NbC addition at low NbC content below 1 at.% may be attributed to the solid solution strengthening due to C atoms soluble in the constituent phases Ni3 Al and/or Ni3 V. It has been reported that boron (B) and carbon (C) atoms behave as interstitial in Ni3 Al [21,26] and result in much significant strengthening [27–29] than the majority of substitutional atoms in Ni3 Al [30–32]. The yield strength increment per
atom fraction normalized by the shear modulus, G of Ni3 Al using the value of 6.5 × 104 MPa was ∼2 G for B and C atoms and ∼1/3 G for most of substitutional atoms in Ni3 Al [27–32]. The further addition of NbC content beyond about 1 at.% did not result in further increase of the flow strength. In the alloys with high NbC content, solid solution strengthening due to C atoms is no longer expected because C atoms are saturated exceeding the solubility limit. In high NbC contents, V atoms are slightly depleted and Nb atoms are slightly surplus in the matrix, consequently resulting in that the yield strength becomes less sensitive to NbC content. On the other hand, the apparent decrease of the tensile strength (UTS) in NbC content range beyond 1.0 at.% may be correlated with the
Fig. 8. SEM fractography (taken at a low magnification) of the base (a–c) and 1.0 NbC (d–f) alloys that were deformed at room temperature (RT) (a, d), 1073 K (b, e) and 1173 K (c, f), respectively.
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Fig. 9. SEM fractography (taken at high magnification) of the base (a–c) and 1.0 NbC (d–f) alloys that were deformed at room temperature (RT) (a, d), 1073 K (b, e) and 1173 K (c, f), respectively.
existence of large size Nb carbide dispersions in the matrix. Such large size Nb carbide dispersions promote a premature fracturing, as shown in Fig. 10. No solid solution strengthening due to C atoms was observed at high temperature (i.e., 1073 K and 1173 K) deformation of the NbC added alloys (Fig. 7). Probably, diffusion related deformation mechanism appears to be operated, and therefore is little affected by C atoms soluble in the matrix. Much detailed studies, solid solution strengthening mechanism due to interstitial atoms as well as substitutional atoms in especially Ni3 V (D022 ) are certainly required to clarify the mechanisms responsible for the strengthening behavior observed in the NbC added alloys. 4.2. Ductilization by NbC addition The dual two-phase intermetallic alloys have been reported to display certain levels of tensile elongation over a wide range of temperatures [12,14,16,17]. Two constituent phases, Ni3 Al and Ni3 V with a relatively simple structure are plastically deformable in a wide temperature range. The deformation of Ni3 Al satisfies the von Mises criterion for the general plasticity of polycrystalline material, activating 1/21 1 0{1 1 1} slip systems [33]. Also, it has been reported that Ni3 V can be deformed by activating 1 1 2{1 1 1} slip system in a wide temperature range [34]. The dual two-phase intermetallic alloys are composed of fine microstructures and have certain orientation relationships between the constituent phases with highly coherent interfacial structures. The numbers of dislocations piled-up against the interface are reduced in such a
microstructure, and consequently the initiation and subsequent propagation of the micro-cracks are suppressed. Furthermore, boron was doped into the dual two-phase intermetallic alloys and has been shown to be effective of suppressing intergranular fracture at least at low temperatures [19–22]. The improvement of tensile elongation by NbC addition to the dual two-phase intermetallic alloy was observed at all temperatures while the enhancement of the flow strength was confined to low temperature (Fig. 7). Let me first discuss the tensile elongation behavior at low temperatures (i.e., RT and 873 K). The effect of NbC addition on the tensile elongation may be attributed to locally intensified plastic deformation around the Nb carbide dispersions. If the interfacial strength between the matrix and Nb carbide dispersion is high enough, it appears that Nb carbide dispersions do not act as the initiation site of micro-crack path. In fact, the fractography in the NbC added alloys exhibited dimple-like fracture patterns, indicating the evidence of significant plastic deformation around the Nb carbide dispersions. As a result, tensile elongation was enhanced and showed highest values at about 1 at.% NbC content where most appropriate size and density of the Nb carbide dispersions for tensile elongation exist. In the alloy with high NbC contents, apparently reduced tensile elongation may be associated with large size Nb carbide dispersions. At high temperature (i.e., 1173 K), it was shown that NbC addition had the effect of improving the tensile elongation through changing fracture mode from intergranular fracture to transgranular fracture mixed with intergranular fracture, as shown in Figs. 8 and 9. At present, reasonable explanation responding to this
Fig. 10. Fractography of the 5.0 NbC alloy in which preferential fracturing takes place at the Nb carbide dispersion itself or the interface between the Nb carbide dispersion and the matrix. Note that a micro-crack is observed in the Nb carbide dispersion deformed at (a) room temperature and (b) 1073 K. A micro-crack is drawn by a red circle. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of the article.)
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result is impossible. In most cases, it is believed that second-phase dispersions on grain boundaries act as stress concentrator and thereby promote intergranular fracturing. Therefore, the secondphase dispersions on grain boundaries are generally harmful for tensile elongation even at high temperature. As far as the authors know, there are no any mechanisms by which the second-phase dispersions enhance grain boundary cohesion, thereby suppressing intergranular fracturing. One idea is that C atoms released from NbC carbides segregate to grain boundaries and enhance the grain boundary cohesion, thereby suppressing the intergranular fracturing. In fact, necessary C atoms were observed in the matrix (i.e., grain interior) (Table 2). It has been reported in Ni3 Al [27] and Ni3 (Si,Ti) [35] with L12 structure that a trace amount of C atoms suppressed intergranular fracturing through segregating on grain boundaries. Here, it should be pointed out that the enhancement of grain boundary cohesion by C atom segregation is effective especially at high temperature where the enhancement of grain boundary cohesion by B atom segregation is ineffective. Definitely, much more studies are required to specify the mechanisms responsible for the peculiar tensile ductility improvement observed in high-temperature deformation of the NbC added alloys. Last, the tensile elongation improvement by grain size refinement cannot be excluded as an additional mechanism. 5. Conclusion Dual two-phase intermetallic alloy that has an alloy composition of Ni75 Al9 V13 Nb3 and is composed of geometrically close packed (GCP) Ni3 Al (L12 ) and Ni3 V (D022 ) phases was charged by NbC powders. The microstructure, room-temperature hardness, and high-temperature tensile properties of NbC added alloys were investigated. The following results were obtained from the present study. (1) The two-phase microstructures defined by primary Ni3 Al precipitates and eutectoid region consisting of Ni3 Al and Ni3 V phases were kept regardless of NbC content and were dispersed by the Nb carbide particles. (2) EPMA results showed that C atoms and slightly surplus Nb atoms stemming from NbC carbides were contained in the matrix of the NbC added alloys. (3) The yield strength as well as tensile strength showed a peak at intermediate temperature, followed by a decrease at high temperature at all NbC added alloys. (4) In low temperature range the yield strength rapidly increased with increasing NbC content up to1 at.% NbC content while in high temperature range the strengthening by NbC addition little took place. (5) The tensile elongation increased with increasing NbC content up to 1 at.% NbC content and was independent of NbC content over 1 at.% NbC content in whole temperature ranges.
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(6) The NbC addition changed fracture mode from transgranular fracture with cleavage-like patterns to that with dimple-like patterns in low temperature range, and from intergranular fracture to ductile transgranular mixed intergranular fracture in high temperature range. (7) The observed strengthening up to about 1.0 at.% NbC content was attributed to the solid solution strengthening by C atoms released from the charged NbC carbides. (8) The observed tensile elongation enhancement was suggested to be due to grain boundary segregation of C atoms released from the charged NbC carbides. Acknowledgement This work was supported in part by Grant-in-aid for Scientific Research (S) for the Ministry of Education, Culture, Sports and Technology. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31] [32] [33] [34] [35]
J.H.J. Van Vucht, Less-Common Met. 11 (1966) 308. A.K. Sinha, Trans. Met. Soc. AIME 245 (1969) 911. A.K. Sinha, Prog. Mater. Sci. Part 2 15 (1972) 79. D.P. Pope, S.S. Ezz, Int. Mater. Rev. 29 (1984) 136. K. Tomihisa, Y. Kaneno, T. Takasugi, Intermetallics 10 (2002) 247. K. Ohira, Y. Kaneno, T. Takasugi, J. Mater. Sci. 39 (2004) 2295. K. Tomihisa, Y. Kaneno, T. Takasugi, Intermetallics 12 (2004) 317. K. Ohira, Y. Kaneno, T. Takasugi, Mater. Sci. Eng. A 399 (2005) 332. M. Fujita, Y. Kaneno, T. Takasugi, Mater. Sci. Eng. A 476 (2008) 112. Y. Nunomura, Y. Kaneno, T. Takasugi, Intermetallics 12 (2004) 389. Y. Nunomura, Y. Kaneno, H. Tsuda, T. Takasugi, Acta Mater. 54 (2006) 851. S. Shibuya, Y. Kaneno, M. Yoshida, T. Takasugi, Acta Mater. 54 (2006) 861. S. Shibuya, Y. Kaneno, H. Tsuda, T. Takasugi, Intermetallics 15 (2007) 338. S. Shibuya, Y. Kaneno, M. Yoshida, T. Shishido, T. Takasugi, Intermetallics 15 (2007) 119. W. Soga, Y. Kaneno, T. Takasugi, Intermetallics 14 (2006) 170. W. Soga, Y. Kaneno, T. Takasugi, Mater. Sci. Eng. A 473 (2008) 180. Y. Kaneno, W. Soga, H. Tsuda, T. Takasugi, J. Mater. Sci. 43 (2008) 748. K. Kawahara, Y. Kaneno, T. Takasugi, Intermetallics 17 (2009) 938. W. Huang, Z. Metallkde 88 (1997) 63. E.K. Storms, The Refractory Carbides, Academic Press, 1967, p. 6. L.J. Heutter, H.H. Stadelmaier, Acta Metall. 6 (1958) 367. N.S. Stoloff, Int. Mater. Rev. 34 (1989) 153. M. Yamaguchi, Y. Umakoshi, Prog. Mater. Sci. 34 (1990) 1. K. Hagihara, T. Tanaka, T. Nakano, Y. Umakoshi, Acta Mater. 53 (2005) 5051. E. Orowan, The Symposium on Internal Stresses in Metals and Alloys, Inst. Metals, London, 1948, p. 451. N. Masahashi, T. Takasugi, O. Izumi, Acta Metall. 36 (1988) (1815). N. Masahashi, T. Takasugi, O. Izumi, Acta Metall. 36 (1988) (1823). S.C. Huang, A.I. Taub, K.-M. Chang, Acta Metall. 32 (1984) (1703). C.L. Briant, S.C. Huang, Metall. Trans. 17A (1986) 2084. S. Ochiai, Y. Mishima, M. Yodogawa, T. Suzuki, Trans. JIM 27 (1986) 32. Y. Mishima, S. Ochiai, M. Yodogawa, T. Suzuki, Trans. JIM 27 (1986) 41. Y. Mishima, S. Ochiai, N. Hamao, M. Yodogawa, T. Suzuki, Trans. JIM 27 (1986) 648. R. von, Mises, Z. Ang. Math. Mech. 8 (1928) 161. K. Hagihara, M. Mori, Y. Umakoshi, MRS Proc. Publ. 980 (2007) II05–20. T. Takasugi, M. Yoshida, J. Mater. Sci. 26 (1991) 3032.
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Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Effect of NbC addition on mechanical properties of dual two-phase Ni3 Al–Ni3 V intermetallic alloy Y. Kitaura, Y. Kaneno, T. Takasugi ∗ Department of Materials Science, Osaka Prefecture University, 1-1 Gakuen-cho, Naka-ku, Sakai, Osaka, 599-8531, Japan
a r t i c l e
i n f o
Article history: Received 4 February 2010 Received in revised form 28 May 2010 Accepted 4 June 2010
Keywords: Multiphase intermetallics Mechanical properties at high temperatures Heat treatment Microstructure NbC carbide
a b s t r a c t Dual two-phase intermetallic alloy that has an alloy composition of Ni75 Al9 V13 Nb3 and is composed of geometrically close packed (GCP) Ni3 Al (L12 ) and Ni3 V (D022 ) phases was studied, focusing on the effect of NbC addition on high-temperature tensile properties. The two-phase microstructures defined by primary Ni3 Al precipitates and eutectoid region consisting of Ni3 Al and Ni3 V phases were kept and were dispersed by Nb2 C particles regardless of NbC content. EPMA analysis showed that alloy compositions in matrix of the NbC added alloy were different from those in the base alloy. The positive temperature dependence of flow strength was observed for all NbC added alloys. In low temperature range the maximum strengthening took place at 1 at.% NbC content. On the other hand, tensile elongation increased with increasing NbC content up to 1 at.% NbC content and gradually decreased at NbC content over 1 at.% content in whole temperature ranges. Also, the NbC addition resulted in change of the fracture mode from brittle transgranular fracture to ductile transgranular fracture in low temperature range, and from brittle intergranular fracture to ductile transgranular in high temperature range. Possible mechanisms responsible for the strengthening and ductilization by NbC addition were discussed, based on the behavior of C solutes in the matrix, which were released from the charged NbC carbides. © 2010 Elsevier B.V. All rights reserved.
1. Introduction Geometrically close packed (GCP) Ni3 X intermetallic phases generally exhibit high phase and microstructural stabilities up to their melting points because of low atomic diffusivity owing to their close packed structures [1–4]. Multi-phase intermetallic alloys based on GCP Ni3 X structures have been studied for various Ni3 X (X: Al, Ti and Nb [5,6], X: Si, Ti and Nb [7–9], X: Al, Ti and V [10–14], and X: Al, Nb and V [15–18]) alloy systems to develop a new type of high-temperature structural materials. Among so far studied alloy systems, so-called dual two-phase intermetallic alloys based on Ni3 Al (L12 )–Ni3 V (D022 ) pseudo-binary alloy system were shown to be attractive specifically as high temperature structural materials. The attractive mechanical properties arise from a peculiar microstructure consisting of upper two-phase microstructure with a micron scale and lower two-phase microstructure with a sub-micron scale. The upper two-phase microstructure is composed of primary Ni3 Al precipitates and Ni solid solution (A1) (channel region) at high temperature, and the prior A1 phase is decomposed into Ni3 Al + Ni3 V by a eutectoid reaction at low temperature, resulting in the lower two-phase microstructure. The dual two-phase microstructures exhibit coherent and fine
∗ Corresponding author. Tel.: +81 72 254 9314; fax: +81 72 254 9912. E-mail address: [email protected] (T. Takasugi). 0921-5093/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2010.06.014
microstructure and high microstructural stability [10,11,13,14]. In addition, high temperature mechanical properties (such as tensile and creep strength) of the dual two-phase intermetallic alloys were found to be superior to those of many conventional super alloys [12,14,18]. Considering the development of the dual two-phase intermetallic alloys as next generation-type high temperature structural materials, it is necessary to furthermore improve their mechanical properties. Along this strategy, traditional strengthening or toughening mechanisms that have been employed to many conventional steels and Ni alloys should be applied. Actually, Ti [12,14] and Nb [16–18] resulting in solid solution strengthening have been so far employed to the dual two-phase intermetallic alloys, and also boron [12,16,17] improving tensile elongation has been doped into the dual two-phase intermetallic alloys. In this study, expecting dispersion strengthening due to Nb carbides and/or solid solution strengthening due to C atoms released from Nb carbides, NbC carbides ranging to 5.0 at.% were charged to the dual two-phase intermetallic alloys composed of Ni3 Al and Ni3 V phases containing Nb and doped with boron. Microstructure, room-temperature hardness, and tensile properties of the dual two-phase intermetallic alloys were investigated as functions of NbC content and test temperature. Improved tensile strength and elongation were observed in certain levels of NbC content and in a wide range of test temperature. Strengthening and ductilization mechanisms due to Nb carbides were discussed, based on the experimental results.
Y. Kitaura et al. / Materials Science and Engineering A 527 (2010) 6012–6019 Table 1 Alloy compositions used in this study (at.%).
Base 0.2 NbC 0.5 NbC 1.0 NbC 2.5 NbC 5.0 NbC a
Ni
Al
V
Nb
NbCa
Ba (wt ppm)
75 75 75 75 75 75
9 9 9 9 9 9
13 13 13 13 13 13
3 3 3 3 3 3
– 0.2 0.5 1.0 2.5 5.0
100 100 100 100 100 100
NbC and B are an extra number ; the numbers are not included in the total.
2. Experimental procedures Alloy buttons with a diameter of 50 mm were prepared by arc melting in argon gas atmosphere using tungsten electrode and copper hearth. The raw materials used in this study were 99.9 wt.% Ni, 99.99 wt.% Al, 99.9 wt.% V, 99.9 wt.% Nb, 99.9 wt.% boron (B) and 99 wt.% NbC. Boron was doped to suppress grain boundary fracture of the present intermetallic alloy [12,16,17]. For NbC carbides, powders with particle sizes of 1–3 m were used. Each alloy button was remelted more than three times to ensure chemical homogeneity through an entire button cross section. A composition expressed by Ni75 Al9 V13 Nb3 (at.%) doped with 100 wt ppm B is a base alloy composition adopted in this study. NbC was charged up to 5.0 at.% to the base alloy composition, as shown in Table 1. Hereafter, these alloy compositions are referred as base, 0.2 NbC, 0.5 NbC, 1.0 NbC, 2.5 NbC, and 5.0 NbC alloys in the followings. These alloys were homogenized at 1553 K for 3 h in a dynamic vacuum condition and then rapidly furnace cooled to room temperature to obtain the dual two-phase microstructure. Vickers hardness test was conducted in conditions of a holding time of 20 s and a load of 1 kg, using the plate-like specimens with a thickness of 1 mm that were sliced from the arc melted buttons. Hardness data points more than ten were collected and averaged in each experimental condition. Microstructures and alloy compositions in matrix and dispersions of the NbC added alloys were investigated by optical microscopy (OM) and scanning electron microscopy (JEOL: JXA870) with a wavelength-dispersive spectroscopy (WDS). The specimens were sliced from the heat-treated arc buttons along a solidified direction using an electro discharge machine (EDM). The sliced specimens were mechanically abraded on SiC paper and then electronically polished in a mixed solution of 15 ml H2 SO4 + 85 ml CH3 OH at 243 K. Also, X-ray diffraction (Rigaku: RINT-2500HKLC) was performed to determine constituent phases and second-phases, using Ni-filtrated Cu K␣ radiation at an accelerated voltage of 30 kV.
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Tensile specimens with a gauge dimension of 2 mm × 1 mm × 10 mm were cut from the heat-treated arc buttons using an EDM. Tensile axis of the tensile specimen was chosen in perpendicular to solidification direction of the arc button. The prepared tensile specimens were composed of bamboo-type grains in which most grain boundary planes are aligned in perpendicular to the tensile axis. Tensile tests were conducted in a temperature range between room temperature and 1173 K within a metal tube surrounded by an electronic furnace. The electro furnace was evacuated to a vacuum degree of approximately 1.5 × 10−3 Pa before testing. The tensile specimens were deformed at a strain rate of 1.66 × 10−4 s−1 up to fracturing. Tensile elongation was calculated from load-elongation curve drawn in the recorded chart, and defined as plastic strain to fracture. The fracture surface of the tensile deformed specimens was examined by using SEM. 3. Results 3.1. Microstructure Large columnar grains that were preferentially elongated along solidification direction and have a grain size of several mm were observed in arc melted material of the base alloy. The addition of NbC exceeding 0.2 at.% resulted in grain refinement and the grain size decreased from several mm to several hundred m with increasing NbC content as clearly shown in Fig. 1. The columnar grains similar to the base alloy were observed in all NbC added alloys. Fig. 2 shows SEM images of microstructures of the alloys observed at a low magnification. The base alloy and 0.2 NbC alloy exhibited a featureless microstructure consisiting of Ni3 Al + Ni3 V although two phases are not resolved in this SEM image. In the NbC added alloys (0.2 NbC, 0.5 NbC, 1.0 NbC and 5.0 NbC alloys are representatively shown in Fig. 2), there were a number of large bright Nb carbides with size of several m within grain interiors. It is found from these pictures that the size and density of the Nb carbides tend to increase with increasing NbC content, and also the Nb carbides tend to coalesce with increasing NbC content. It is suggested that the Nb carbides are formed as a primary phase and present in dendritic core region during solidification. X-ray diffraction patterns of base, 0.2 NbC, 0.5 NbC, 1.0 NbC, and 5.0 NbC alloys are shown in Fig. 3. In the base alloy, only diffraction lines indexed by both Ni3 Al and Ni3 V were identified. In the other NbC added alloys, additional diffraction lines indexed by NbC were identified and their lines were counted at higher angles than those calculated from a stoichiometric NbC [19,20]. This result reveals that the Nb carbides consisting of low C content, in other words, the Nb carbides to be assumed as the Nb2 C carbides are contained in
Fig. 1. Cross-section views of arc melted buttons of the (a) base, (b) 0.2 NbC, (c) 1.0 NbC, and (d) 5.0 NbC alloys.
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Fig. 2. SEM images of microstructures of the (a) base, (b) 0.2 NbC, (c) 0.5 NbC, (d) 1.0 NbC, and (e) 5.0 NbC alloys. Pictures were taken at a low magnification to show grain size and morphology, and NbC dispersions.
Fig. 3. X-ray diffraction patterns of the (a) base, (b) 0.2 NbC, (c) 0.5 NbC, (d) 1.0 NbC, and (e) 5.0 NbC alloys.
the dual two-phase intermetallic alloys, consistent with SEM-EPMA result shown later. SEM-EPMA results in which alloy compositions in matrix and Nb carbides of the base and 5.0 NbC alloys are shown in Table 2. Here, it is noted that the spot size for the EPMA observation for the matrix was large enough to examine average composition of the region decomposed into Ni3 Al and Ni3 V while that for the carbides was reduced as small as possible. Also for the analysis of carbon content, the standard specimen was used for obtaining accurate results. Alloy composition in the base alloy was almost equal to the nominal alloy composition, i.e., the intended alloy composition Table 2 Alloy compositions (at.%) in matrix and carbide in the base and 5.0 NbC alloys. Base
Ni
Al
V
Nb
Matrix
74.8
8.6
12.3
3.4
C –
5.0 NbC
Ni
Al
V
Nb
C
Matrix Carbide
74.1 2.43
9.4 0.016
11.3 8.66
3.8 60.9
1.42 28.0
shown in Table 1. In the Nb carbides in the 5.0 NbC added alloy, C content is much reduced from NbC. According to the composition of the Nb carbide it seems to be not NbC but Nb2 C. However, detailed analysis was not conducted in the crystal structure of the Nb carbides. Also, it was found that V atoms which are constituent element in the base alloy are contained in the Nb carbides. On the other hand, C atoms and surplus Nb atoms by about 0.8 at.%, which are assumed to be released from NbC carbides, are contained in the matrix in the 5.0 NbC added alloy. SEM images of microstructures in grain interior (matrix) of the alloys observed at a high magnification are shown in Fig. 4. The upper two-phase microstructures consisting of primary Ni3 Al precipitates and channel region is clearly imaged in each alloy. The size and morphology of the upper two-phase microstructures are similar and insensitive to NbC content. For the lower two-phase microstructure, i.e., the eutectoid region (channel region) corresponding to the base alloy composition, an extensive study has been done [16,17]. The eutectoid region, i.e. the prior Ni solid solution (A1) was shown to be composed of lamellar-like structures consisting of Ni3 Al and Ni3 V.
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Fig. 4. SEM images of microstructures of the (a) base, (b) 0.2 NbC, (c) 1.0 NbC, and (d) 5.0 NbC alloys. Pictures were taken at a high magnification to show the dual two-phase microstructure of each alloy.
3.2. Mechanical properties Change of room-temperature Vickers hardness with NbC content is shown in Fig. 5. As NbC content increases, the hardness monotonously increases up to about 2 at.% NbC content and then appears to be saturated in NbC content range beyond 2 at.%. Yield strength, tensile strength (UTS) and tensile elongation for the base, 0.2 NbC, 0.5 NbC, 1.0 NbC, 2.5 NbC, and 5.0 NbC alloys are shown in Fig. 6, respectively, as a function of temperature. In all alloys, positive temperature dependence of the yield strength took place, and at a peak temperature (around 900 K), the tensile strength as well as the yield strength almost exceeded 1 GPa for all alloys. The flow strength increase by NbC addition generally occurred in low temperature range below a peak temperature but little occurred in high temperature range beyond a peak temperature. The details will be described later. On the other hand, the improvement of tensile elongation by NbC addition was observed for all test temperatures. The details also will be described later.
Fig. 5. Change of Vickers hardness by NbC content.
Variations of yield strength, tensile strength (UTS), and tensile elongation with NbC content at room temperature (RT), 873 K, 1073 K, and 1173 K are shown in Fig. 7, respectively. It was clearly observed that in low temperature range (e.g. RT and 873 K) below a peak temperature, the yield strength as well as the tensile strength rapidly increased with increasing NbC content and made a peak at around 1.0 at.% NbC content, followed by a steady decrease (in the case of the tensile strength) or by consistently high level (in the case of the yield strength) in high NbC content range. The yield strength increase by about 450 MPa was thus observed at room temperature by the addition of 1 at.% NbC (Fig. 7). In high temperature range (e.g. 1073 K and 1173 K) beyond a peak temperature, the flow strength was little affected by NbC addition in their whole contents. Looking the variation curves of tensile elongation vs. NbC content, the tensile elongation increased with increasing NbC content and made a peak at around 1.0 at.% NbC content, followed by a decrease in high NbC content range. In the base alloy, preferential fracture mode at low temperature (below a peak temperature) was transgranular fracture accompanied by cleavage-like pattern while that at high temperature (beyond a peak temperature) was intergranular fracture mode accompanied by grain boundary facets on which tinny dimple patterns were observed, as shown in Figs. 8 and 9. However, fracture mode was apparently changed by NbC addition. That is, ductile transgranular fracture mode accompanied by higher fraction of dimple patterns was prevailing the fracture pattern in the entire deformation temperature tested (not only low temperature but also high temperature), as shown in Figs. 8 and 9. The observed fracture patterns reveal that largely energy-dissipated event occurs during propagation of micro-crack in the NbC added alloys not only at low temperature but also at high temperature, consequently improving the tensile elongation. Thus, the tensile behavior is consistent with the observed fractography: once grain boundary fracturing is suppressed, high tensile elongation which the dual two-phase microstructure intrinsically has can be obtained, resulting in the improvement of tensile ductility. However, large Nb carbides may act as the initiation sites of fracturing and may reduce tensile ductility of the present alloys. Fig. 10 shows the fractography of the
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Fig. 6. Yield strength, tensile strength (UTS), and tensile elongation as a function of test temperature for the (a) base, (b) 0.2 NbC, (c) 0.5 NbC, (d) 1.0 NbC, (e) 2.5 NbC, and (f) 5.0 NbC alloys, respectively.
5.0 at.% NbC added alloy in which preferential fracturing takes place from the Nb carbides themselves or the interface between Nb carbide and the matrix.
clarify precipitation phenomenon of the Nb carbides in solid state is an important issue to more sophistically control the dispersion of the Nb carbide particles and to improve mechanical properties of the dual two-phase intermetallic alloys.
4. Discussion 4.1. Strengthening by NbC addition It is assumed that the charged NbC carbides were wholly melted into liquid phase during arc melting and redeposited during solidification process although NbC with a cubic structure has very high formation enthalpy energy and therefore a very high melting point (3898 K) [19,20]. The result that the Nb carbides analyzed by EPMA showed different chemical composition from the charged NbC carbides, and at the same time surplus Nb atoms and C atoms were detected from the matrix of the NbC added alloys (Table 2) supports this speculation. Also, the morphology and size of the Nb carbides observed in the microstructures were apparently different from those of the charged NbC carbides. It is also speculated from SEM-EPMA (Fig. 2 and Table 2) and XRD (Fig. 3) results that the addition of NbC contents exceeding 0.5–1 at.% is high enough resulting in C solute atoms beyond the solubility limit (∼2 at.%) in the matrix of the present alloy. This value is not so much different from the reported solubility limit (several at.%) of C atoms in the major constituent phase Ni3 Al in the present alloy (composition) [21]. However, the lattice expansion due to C atoms soluble in Ni3 Al and/or Ni3 V was not correctly evaluated from X-ray diffraction patterns because the majorities of diffraction lines from both phases were superimposed and therefore are ambiguous and too broad. To
The strength properties observed in the base alloy basically are interpreted by the interfacial strengthening in addition to the strengthening due to two constituent phases [18]. The latter strengthening was basically interpreted in terms of the mixture rule of strength of two constituent phases [18]. The yield strength anomaly shown in Fig. 6 primarily arises from the nature that Ni3 Al and Ni3 V display the positive temperature dependence of the flow strength [4,22–24]. In both phases, the anisotropy in anti-phase boundary (APB) energy between the glide plane and non-glide plane promotes thermally activated cross-slip of the dissociated screw dislocation from the glide plane to non-glide plane, resulting in the strength anomaly. Besides the strength anomaly, large interfacial strengthening occurs in the present dual two-phase microstructures when a dislocation passes through the interface between the Ni3 Al and Ni3 V phases or between different variants in the Ni3 V phase [18]. Regarding mechanisms responsible for the additional strengthening due to NbC addition, the strengthening mechanism based on grain refinement is at least unlikely to be applied because the observed flow strength did not monotonously increase as
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Fig. 7. Variation of yield strength, tensile strength (UTS), and tensile elongation with NbC content at (a) room temperature (RT), (b) 873 K, (c) 1073 K, and (d) 1173 K, respectively.
grain size decreases (i.e., as NbC content increases). Also, the dispersion strengthening due to the Nb carbides is not applicable. A qualitative evaluation assuming the Orowan-type mechanism [25] predicts flow strength increase of several 10 MPa, apparently much lower than the strength increase (∼450 MPa) observed in the present study. Instead, the strengthening by NbC addition at low NbC content below 1 at.% may be attributed to the solid solution strengthening due to C atoms soluble in the constituent phases Ni3 Al and/or Ni3 V. It has been reported that boron (B) and carbon (C) atoms behave as interstitial in Ni3 Al [21,26] and result in much significant strengthening [27–29] than the majority of substitutional atoms in Ni3 Al [30–32]. The yield strength increment per
atom fraction normalized by the shear modulus, G of Ni3 Al using the value of 6.5 × 104 MPa was ∼2 G for B and C atoms and ∼1/3 G for most of substitutional atoms in Ni3 Al [27–32]. The further addition of NbC content beyond about 1 at.% did not result in further increase of the flow strength. In the alloys with high NbC content, solid solution strengthening due to C atoms is no longer expected because C atoms are saturated exceeding the solubility limit. In high NbC contents, V atoms are slightly depleted and Nb atoms are slightly surplus in the matrix, consequently resulting in that the yield strength becomes less sensitive to NbC content. On the other hand, the apparent decrease of the tensile strength (UTS) in NbC content range beyond 1.0 at.% may be correlated with the
Fig. 8. SEM fractography (taken at a low magnification) of the base (a–c) and 1.0 NbC (d–f) alloys that were deformed at room temperature (RT) (a, d), 1073 K (b, e) and 1173 K (c, f), respectively.
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Fig. 9. SEM fractography (taken at high magnification) of the base (a–c) and 1.0 NbC (d–f) alloys that were deformed at room temperature (RT) (a, d), 1073 K (b, e) and 1173 K (c, f), respectively.
existence of large size Nb carbide dispersions in the matrix. Such large size Nb carbide dispersions promote a premature fracturing, as shown in Fig. 10. No solid solution strengthening due to C atoms was observed at high temperature (i.e., 1073 K and 1173 K) deformation of the NbC added alloys (Fig. 7). Probably, diffusion related deformation mechanism appears to be operated, and therefore is little affected by C atoms soluble in the matrix. Much detailed studies, solid solution strengthening mechanism due to interstitial atoms as well as substitutional atoms in especially Ni3 V (D022 ) are certainly required to clarify the mechanisms responsible for the strengthening behavior observed in the NbC added alloys. 4.2. Ductilization by NbC addition The dual two-phase intermetallic alloys have been reported to display certain levels of tensile elongation over a wide range of temperatures [12,14,16,17]. Two constituent phases, Ni3 Al and Ni3 V with a relatively simple structure are plastically deformable in a wide temperature range. The deformation of Ni3 Al satisfies the von Mises criterion for the general plasticity of polycrystalline material, activating 1/21 1 0{1 1 1} slip systems [33]. Also, it has been reported that Ni3 V can be deformed by activating 1 1 2{1 1 1} slip system in a wide temperature range [34]. The dual two-phase intermetallic alloys are composed of fine microstructures and have certain orientation relationships between the constituent phases with highly coherent interfacial structures. The numbers of dislocations piled-up against the interface are reduced in such a
microstructure, and consequently the initiation and subsequent propagation of the micro-cracks are suppressed. Furthermore, boron was doped into the dual two-phase intermetallic alloys and has been shown to be effective of suppressing intergranular fracture at least at low temperatures [19–22]. The improvement of tensile elongation by NbC addition to the dual two-phase intermetallic alloy was observed at all temperatures while the enhancement of the flow strength was confined to low temperature (Fig. 7). Let me first discuss the tensile elongation behavior at low temperatures (i.e., RT and 873 K). The effect of NbC addition on the tensile elongation may be attributed to locally intensified plastic deformation around the Nb carbide dispersions. If the interfacial strength between the matrix and Nb carbide dispersion is high enough, it appears that Nb carbide dispersions do not act as the initiation site of micro-crack path. In fact, the fractography in the NbC added alloys exhibited dimple-like fracture patterns, indicating the evidence of significant plastic deformation around the Nb carbide dispersions. As a result, tensile elongation was enhanced and showed highest values at about 1 at.% NbC content where most appropriate size and density of the Nb carbide dispersions for tensile elongation exist. In the alloy with high NbC contents, apparently reduced tensile elongation may be associated with large size Nb carbide dispersions. At high temperature (i.e., 1173 K), it was shown that NbC addition had the effect of improving the tensile elongation through changing fracture mode from intergranular fracture to transgranular fracture mixed with intergranular fracture, as shown in Figs. 8 and 9. At present, reasonable explanation responding to this
Fig. 10. Fractography of the 5.0 NbC alloy in which preferential fracturing takes place at the Nb carbide dispersion itself or the interface between the Nb carbide dispersion and the matrix. Note that a micro-crack is observed in the Nb carbide dispersion deformed at (a) room temperature and (b) 1073 K. A micro-crack is drawn by a red circle. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of the article.)
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result is impossible. In most cases, it is believed that second-phase dispersions on grain boundaries act as stress concentrator and thereby promote intergranular fracturing. Therefore, the secondphase dispersions on grain boundaries are generally harmful for tensile elongation even at high temperature. As far as the authors know, there are no any mechanisms by which the second-phase dispersions enhance grain boundary cohesion, thereby suppressing intergranular fracturing. One idea is that C atoms released from NbC carbides segregate to grain boundaries and enhance the grain boundary cohesion, thereby suppressing the intergranular fracturing. In fact, necessary C atoms were observed in the matrix (i.e., grain interior) (Table 2). It has been reported in Ni3 Al [27] and Ni3 (Si,Ti) [35] with L12 structure that a trace amount of C atoms suppressed intergranular fracturing through segregating on grain boundaries. Here, it should be pointed out that the enhancement of grain boundary cohesion by C atom segregation is effective especially at high temperature where the enhancement of grain boundary cohesion by B atom segregation is ineffective. Definitely, much more studies are required to specify the mechanisms responsible for the peculiar tensile ductility improvement observed in high-temperature deformation of the NbC added alloys. Last, the tensile elongation improvement by grain size refinement cannot be excluded as an additional mechanism. 5. Conclusion Dual two-phase intermetallic alloy that has an alloy composition of Ni75 Al9 V13 Nb3 and is composed of geometrically close packed (GCP) Ni3 Al (L12 ) and Ni3 V (D022 ) phases was charged by NbC powders. The microstructure, room-temperature hardness, and high-temperature tensile properties of NbC added alloys were investigated. The following results were obtained from the present study. (1) The two-phase microstructures defined by primary Ni3 Al precipitates and eutectoid region consisting of Ni3 Al and Ni3 V phases were kept regardless of NbC content and were dispersed by the Nb carbide particles. (2) EPMA results showed that C atoms and slightly surplus Nb atoms stemming from NbC carbides were contained in the matrix of the NbC added alloys. (3) The yield strength as well as tensile strength showed a peak at intermediate temperature, followed by a decrease at high temperature at all NbC added alloys. (4) In low temperature range the yield strength rapidly increased with increasing NbC content up to1 at.% NbC content while in high temperature range the strengthening by NbC addition little took place. (5) The tensile elongation increased with increasing NbC content up to 1 at.% NbC content and was independent of NbC content over 1 at.% NbC content in whole temperature ranges.
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(6) The NbC addition changed fracture mode from transgranular fracture with cleavage-like patterns to that with dimple-like patterns in low temperature range, and from intergranular fracture to ductile transgranular mixed intergranular fracture in high temperature range. (7) The observed strengthening up to about 1.0 at.% NbC content was attributed to the solid solution strengthening by C atoms released from the charged NbC carbides. (8) The observed tensile elongation enhancement was suggested to be due to grain boundary segregation of C atoms released from the charged NbC carbides. Acknowledgement This work was supported in part by Grant-in-aid for Scientific Research (S) for the Ministry of Education, Culture, Sports and Technology. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31] [32] [33] [34] [35]
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