Since Aoki and Izumi found out the effectiveness of B-doping for improving the ductility of Ni3Al alloys [1], the ductilizing and toughening mechanisms of the doping have been actively discussed [2±5]. In the early stage of the studies, the segregation of B on the grain boundaries was found to have a bene®cial effect on enhancing the grain boundary cohesion and to relieve the intrinsic grain boundary brittleness [2, 3]. Recently, the suppression of environmental embrittlement by the doping is believed to be the main factor for the toughening of Ni3Al alloys [4]. Because of the diffusivity of atomic hydrogen, which, when generated by the surface reaction of moisture in air, is reduced [5], the alloys are consequently insensitive to the environmental effect by the doping. However, there are some problems that cannot be explained except by the suppression of environmental effects such as the alloy composition dependence of the ef®cacy of B [2]. In the present work, the toughening mechanism of stoichiometric and Ni-rich Ni3Al alloys by the doping will be brie y discussed. Ni3Al alloys, with and without the doping of 0.1 at % B, were fabricated by reactive hot-pressing [6] at 1573 K with an applied pressure of 30 MPa. The fracture toughness of the alloys was measured by the single-edge chevron notched beam (SECNB) method in air and an oil bath (silicone oil SRX310: Toray Dow Corning Silicone Co., Ltd) at 300 K. The chevron notched (notch length a0: 4 mm) testing bars with 8 mm 3 4 mm 3 35 mm dimensions were used for the evaluation of the fracture toughness. The three-point bending test of the notched bars with a span length of 30 mm was conducted at a loading rate of 10y2 y 10 MPa m secy1. The tensile test of the alloys was also performed in air by using the specimens with a gage section of 0:5 mm 3 1 mm 3 8 mm. The strain rate for the tensile test was 8:3 3 10y4 secy1. The fracture surface of the specimens was observed by scanning electron microscopy (SEM). Fig. 1 shows the loading rate dependence of fracture toughness for the Ni3Al alloys measured by SECNB method. The fracture toughness of the alloys without the doping decreases with the decrease of the loading rate, although a similar behavior is not observed in the oil bath. Because the moistureinduced embrittlement is expected to be suppressed in the oil bath, this loading rate dependence in air is caused by the environmental effect. However, the fracture toughness measured at a loading rate of 10 MPa m secy1 in air is equal to that measured in the oil bath. Therefore, the fracture toughness unaffected by the environmental effect (intrinsic toughness) can be estimated above 10 MPa m secy1. The intrinsic toughness values for the nondoped 24 and 25 at % Al alloys are estimated to be 19 and 16 MPa m, respectively. The environmental embrittlement of Ni-25 at % Al alloy is successfully inhibited by the 0.1 at % Bdoping because the fracture toughness of the doped alloy in air is independent of the loading rate and is equal to that in the oil bath. Concurrently with the suppression of the environmental effect, its intrinsic fracture toughness is improved to be 30% larger than that of the non-doped alloy by the doping. On the other hand, the chevron-notched testing bars of the B-doped Ni-24 at % Al alloy cannot be fractured validly in the present conditions due to its extensive ductility. At least, the toughness of the doped Ni24 at % Al alloy is much larger than the intrinsic value of the non-doped one. If the doped B only plays a role for the suppression of environmental effect, the fracture toughness of the B-doped Ni24 at % Al alloy is improved until it has the intrinsic toughness of the non-doped alloys. However, the fracture toughness of the B-doped alloy is clearly improved much more. Hence, another effect of B is necessary to interpret the extensive improvement of
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Oct 23, 2010
Z.W Zhang + 10
Open Access