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Synthesis and properties of bulk metallic glasses in the ternary Ni–Nb–Zr alloy system
Materials Science and Engineering A 492 (2008) 221–229
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Synthesis and properties of bulk metallic glasses in the ternary Ni–Nb–Zr alloy system Z.W. Zhu a,b , H.F. Zhang a,∗ , B.Z. Ding a , Z.Q. Hu a a Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang 110016, China b Graduate School of the Chinese Academy of Sciences, Beijing 100039, China
a r t i c l e
i n f o
Article history: Received 6 November 2007 Received in revised form 11 March 2008 Accepted 8 April 2008 Keywords: Bulk metallic glass Ni-based alloy Thermal property Mechanical property Corrosion resistance
a b s t r a c t Bulk metallic glasses (BMGs) with high thermal stability, good mechanical properties and high corrosion resistance were synthesized in the Ni–Nb–Zr system. A large bulk glass-forming region with 60 < Ni < 64, 28 < Nb < 38 and 0 < Zr < 9 (in at.%) was found. The critical size for the glass formation is 3 mm. These investigated Ni-based BMGs process high glass transition temperature of about 880–900 K and high onset crystallization temperature of 915–932 K as well as high compressive fracture strength of approximate 3.0–3.2 GPa along with some compressive plasticity of about 2%. Electrochemical measurements indicate they also exhibit high corrosion resistance, i.e., large passive region above 1.5 V (vs. saturated calomel reference electrode, SCE). The influence of the Zr content on the glass-forming ability (GFA) and corrosion behaviors was carefully studied, indicating that some Zr addition improves the GFA and corrosion resistance. © 2008 Elsevier B.V. All rights reserved.
1. Introduction Bulk metallic glasses (BMGs, typically referred to a minimum casting dimension larger than 1 mm) have been greatly concerned in the past few decades because they are of particular scientific and engineering interests [1–3]. Some progress in both glass-forming ability (GFA) and mechanical properties has been made, for example, the amorphous samples with the critical size over 10 mm were successfully prepared in Mg [4], Zr [5], Fe [6], Ti [7], Cu [8], Pd [9], etc., based alloys, and some Cu [10,11], Zr [12,13], Ti [7,14] based BMGs samples display very large compressive plastic strain in compression tests. It is exciting even though some problems are still puzzled. Meanwhile, due to various desirable properties, including high yield strength, hardness and elastic strain limit in addition to reasonably high fracture toughness, fatigue resistance and corrosion resistance, etc., BMGs have been tried to be made into some items such as sporting goods, surgical instruments, and strong, thin cases for electronic devices such as mobile phone and U-disc [3]. To satisfy the requirements of commercial applications, it is urgent to improve the known BMGs’ plasticity or to develop new BMGs with higher GFA and better mechanical properties, especially based on common metals, such as Al, Cu, Fe, Ni, etc.
∗ Corresponding author. Fax: +86 2423971783. E-mail address: [email protected] (H.F. Zhang). 0921-5093/$ – see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2008.04.021
In the case of Ni-based alloys, bulk metallic glasses were prepared in the complex alloy systems, such as Ni–Nb–Cr–Mo–P–B [15], Ni–Ti–Zr–(Si,Sn) [16], Ni–Nb–Ti–Zr–Co–Cu [17], Ni–Nb–Ti–Zr–Si–Sn [18,19], Ni–Nb–Sn [20,21], Ni–Cu–Ti–Zr–Al [22], Ni–Ta–Sn [23], etc. However, with compared to that of Cu-, Zr-, Ti- and Fe-, etc., based BMGs [4–9], the GFA of Ni-based amorphous alloys is a challenging subject. Because so far the maximum dimension of Ni-based BMGs’ samples is only 5 mm [22,24], the development of new Ni-based glass former with higher GFA is imperative. In the period of the conventional amorphous alloys, Ni–Nb and Ni–Zr systems were famous for their GFA. Very recently, BMG samples up to 2-mm thick were fabricated in binary Ni–Nb alloy system [25,26]. Some reported works also indicate that Ni-based Nb-bearing BMGs possess better mechanical properties than other Ni-based ones [16,27]. According to Miracle’s efficient cluster packing model [28], Ni–Nb–Zr system with good atomic size distribution (the Goldschmidt atomic radius of Zr is 0.160 nm, which is larger than 0.146 nm and 0.128 nm for Nb and Ni, respectively) could have high GFA [29,30]. Additionally, some believe that some valve metals, such as Nb and Zr, etc., enrich in the surface film to prevent the materials from corrosion [31–34]. As a result, the ternary Ni–Nb–Zr system might be a good candidate to develop new Ni-based BMGs with better combination among GFA, good mechanical properties and high corrosion resistance. Nevertheless, systematic investigation of BMGs in ternary Ni–Nb–Zr system is hardly reported [29,34,35].
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In this paper, we reported that bulk glasses can be formed in quite wide composition range of ternary Ni–Nb–Zr system. The 3mm diameter glassy samples were successfully prepared by copper mold injection casting method. They displayed good mechanical properties, high thermal stability and high anti-corrosion property. The effect of the Zr concentration on the phase transformation upon casting and corrosion behavior is also carefully discussed. 2. Experimental Master alloy ingots were prepared by arc melting a mixture of ultrasonically cleansed Ni, Nb and Zr with a purity of above 99.9% on a water-cooled copper hearth under Ti-gettered high purity argon atmosphere. The chemical homogeneity was obtained by repeated melting at least four times. The ingots were then remelted under high vacuum in a quartz tube by using induction heating coil and injected through a nozzle with 0.5–1 mm in diameter into the copper mould with a cavity of 2–4 mm diameter. The as-cast samples were characterized with X-ray diffraction (XRD, Philips PW1050, Cu K␣), transmission electron microscopy (TEM, JEOL 2010, 200 kV) and differential scanning calorimetry (DSC; Netzsch DSC 404C). The specimens used for XRD measurement were cut from the middle part of the as-cast rods. Thin slices from the 2.5-mm diameter as-cast rods were used for preparing the TEM samples, which were ground and mechanically dimpled with a GATAN precision dimple grinder as well as polished using argon ion milling as the final thinning process using a GATAN precision ion polishing system (PIPS). DSC measurements were performed in a flowing argon atmosphere at a heating rate of 0.33 K/s. Mechanical properties were measured with the samples of 2mm diameter and 4-mm length on a servo-hydraulic materials testing system (MTS 810). To perform compression tests under a constant strain rate of 2 × 10−4 s−1 , a MTS strain gauge was used. Fracture surface was examined by scanning electron microscopy (SEM, Hitachi S3400N). Corrosion behavior of Ni-based BMGs in 1 M HCl aqueous solution open to air was studied by electrochemical measurements on an Advanced Electrochemical System (Princeton Applied Research PARSTAT 2273) at 300 K. Prior to the corrosion tests, the specimens were mechanically polished in cyclohexane with silicon carbide paper up to No. 2000, degreased in acetone, washed in the distilled water, and dried in air. Electrochemical measurements were conducted in a three-electrode cell a platinum counter electrode and a saturated calomel reference electrode (SCE). Potentiodynamic polarization curves were measured at a potential sweep rate of 0.333 mV/s after immersing the samples for several minutes, when the open-circuit potential became almost steady. The surface of the samples exposed to air after mechanical polishing and conducted by potentiodynamic polarization measurements up to 1 V (vs. SCE) in 1 M HCl solution was examined by X-ray photoelectron spectroscopy (XPS) using a photoelectron spectrometer with Al K␣ radiation (h = 1486.6 eV). From these spectra, the composition of the passive film and the underlying alloy surface was quantitatively determined. 3. Results and discussion 3.1. Glass-forming ability The Ni–Nb–Zr alloy system shows good GFA. Fig. 1 illustrates XRD patterns of the as-cast Ni–Nb–Zr alloys rods with a diameter of 2 mm. In Fig. 1a, when a increase from 1 to 9, the 2-mm diameter as-cast samples of Ni61.5 Nb38.5−a Zra alloys display only a series of diffuse maxima around 2 = 42◦ , while for a = 11 the 2-mm diame-
Fig. 1. XRD patterns of the as-cast rods with a diameter of 2 mm for (a) Ni61.5 Nb38.5−a Zra (a = 1, 3, . . ., 11 at.%) [29] and (b) Ni100−b (Nb0.85 Zr0.15 )b (b = 35, 36, . . ., 40 at.%).
ter as-cast rods exhibit apparent crystalline Bragg peaks. It indicates that when Zr content is below 9 at.%, the glassy samples with 2 mm in diameter can be synthesized in the alloys Ni61.5 Nb38.5−a Zra [29]. Likewise, it is also found out that the glass can be also formed in a quite large composition range as the Ni concentration is changed. It is shown in Fig. 1b. When Ni content increases from 60 at.% to 65 at.%, the patterns of the 61–64 at.% Ni-bearing samples consist of only a broad peak without any observable crystalline diffraction peaks, indicating that the 2-mm diameter glassy sample can be made. Outside of this range, crystallization occurs to the samples with 2 mm in diameter. Identification of the crystalline phases will be discussed in detail in Section 3.3. Through tens of alloys experiments, it is discovered that there exists a wide BMG forming region of 60 < Ni < 64, 28 < Nb < 38 and 0 < Zr < 9, in at.%, in ternary Ni–Nb–Zr alloy system, as shown in Fig. 2. In the larger green ellipse region, the samples with at least 2 mm in diameter can be manufactured. It is necessary to point out that the 3-mm diameter as-cast rods are capable to be produced in the smaller purple ellipse region. XRD patterns of the 3-mm diameter samples are shown in Fig. 3, displaying the typical characteristics of those of the amorphous phase. 3.2. Thermal property Characterization of the thermal property of the investigated Ni–Nb–Zr glassy alloys, especially determination of the onset of glass transition temperature (Tg ) and the onset of crystallization
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Fig. 2. Sketch of the glass formation and variation of Tg and Tx in the Ni–Nb–Zr system. The solid green ellipse corresponds to the region where BMGs with at least 2 mm in diameter can be formed, the solid purple ellipse corresponding to 3-mm diameter BMGs.
Fig. 3. XRD patterns of the as-cast rods with a diameter of 3 mm in the three alloys.
temperature (Tx ) as well as liquidus temperature (Tl ) was conducted by DSC measurements. Fig. 4 shows the high temperature DSC profiles recorded at a heating rate of 0.33 K/s. And Table 1 lists the concerning thermal data of the typical alloys. The samples used for DSC measurements were cut from the middle section of the 2-mm diameter as-cast rods. When DSC traces and the thermal properties are considered, the alloys with the Zr content of
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1–9 at.% or Ni content of 61–64 at.% show the distinct glass transition, while for 60 at.% Ni and 65 at.% Ni alloys, no endothermic event can be observed on the traces, implying no glass transition occurs, as shown in Fig. 4a and b. Tg , defined as the onset of the endothermic event, is relatively high and above 880 K. Compared with the known Ni-based BMGs [15–22], it can be inferred that these BMGs have high thermal stability. It is easily found out that for a series of alloys with the composition of Ni61.5 Nb38.5−a Zra , Tg decreases abruptly by about 9 K when the concentration of Zr increases from 3 at.% to 5 at.% (Fig. 4a and Table 1). When Zr content is below 3 at.% or above 5 at.%, Tg nearly remains equal, around 893 K or 882 K, respectively. The similar phenomena on glass transition does not occur to Ni100−b (Nb0.85 Zr0.15 )b alloys. Fig. 4b and Table 1 show that Tg falls continuously as Ni content decreases. Additionally, it is observed from Fig. 4a and b that there exist significant differences in crystallization behaviors. The alloys transform from the beginning three-stage crystallization to double-stage one as the Zr content or Ni content rises. Whether the Zr content increases or Ni content falls, Tx always declines. The variation of Tg and Tx is roughly drawn in Fig. 3. Further, it is found that the variation of Tg is related to the crystallization behavior for the Ni61.5 Nb38.5−a Zra alloys, since the abrupt decrease in Tg occurs at the transition from the three-stage crystallization to double-stage one. As shown in Figs. 2 and 3, the alloy with the 5 at.% Zr exhibited the best GFA of the Ni61.5 Nb38.5−a Zra alloys. As a result, the GFA is inferred to have a close relationship with the crystallization process, which will be discussed in detail in Section 3.3. It is thought that the atomic arrangement configuration is attributed to Tg or Tx dependence of Zr content. As known, the different interaction exists among Ni, Nb and Zr atoms, indicated by different mixing enthalpy values among them, i.e., −49 kJ/mol for Ni–Zr, −30 kJ/mol for Ni–Nb, and 4 kJ/mol for Nb–Zr [36]. Due to the different interaction, it leads to atomic reconfiguration to introduce Zr atoms into the Ni–Nb alloy. Extended X-ray absorption fine structure experiments [35] reveal that the bonds like Ni–Ni and Ni–Zr around Ni atoms and Nb–Ni and Nb–Nb around Nb atoms are chemically preferred to be formed as Zr is added into Ni–Nb alloys. Difference in atomic configuration contributes to the different behaviors, including different Tg , Tx , etc., during reheating process. It also affects the subsequent crystallization as mentioned above. Fig. 4c and d exemplify the melting behaviors of the Ni–Nb–Zr alloys. The liquidus temperature, Tl , and the melting temperature, Tm , decrease with increasing the Zr amount or reducing the Ni amount. The extent of the decline of Tl is faster than that of Tm . Similarly, Trg [37] deduced from the thermal parameters, are proposed to be correlated well with the GFA. They are also given in Table 1. Trg exhibits high values, but does not either possess a good correlation with the GFA in the Ni–Nb–Zr system. The supercooled liquid region T = Tx − Tg , which reflects the thermal stability of the
Table 1 Thermal properties of the as-cast samples with a diameter of 2 mm, except that the diameter of samples in Ni61.5 Nb38.5 alloy is 1.5 mm, deduced from the high temperature DSC measurement at a heating rate of 0.33 K/s Alloy (at.%)
Tg (K)
Tx (K)
Tp (K)
Tm (K)
Tl (K)
Tx (=Tx − Tg , K)
Trg (=Tg /Tl )
Ni61.5 Nb38.5−a Zra
a=0 a=1 a=3 a=5 a=7 a=9
894 893 892 883 882 882
932 926 921 918 913 913
– 935 928 926 923 918
1455 1441 1426 1414 1390 1384
1519 1510 1456 1444 1439 1420
38 33 29 35 31 31
0.589 0.591 0.613 0.612 0.613 0.620
Ni100−b (Nb0.85 Zr0.15 )b
b = 36 b = 37 b = 38 b = 39
899 890 886 884
935 933 924 915
– – – –
1408 1412 1416 1420
1448 1450 1487 1513
36 43 38 31
0.621 0.616 0.599 0.584
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Fig. 4. DSC scans corresponding to glass transitions and crystallizations, melting behaviors of the as-cast rods with a diameter of 2 mm, (a) and (c) for Ni61.5 Nb38.5−a Zra (a = 1, 3, . . ., 9 at.%) [29], (b) and (d) for Ni100−b (Nb0.85 Zr0.15 )b (b = 35, 36, . . ., 40 at.%).
supercooled liquid towards crystallization, varies slightly with the composition (Table 1) and ranges from 30 K to 40 K. It is not directly related to the GFA (shown in Fig. 3) in the current work although T was suggested to characterize the GFA [2]. 3.3. Phase transformation dependence of the Zr content upon solidification Upon solidification, the glass is formed by competing against the nucleation and growth of the crystals in the undercooled melt. The glass can be fabricated under the condition that the nucleation and growth of the primary competing crystals are completely suppressed while the melt is cooled through the temperature interval from Tl (below which crystallization is thermodynamically possible) to Tg (below which the melt is frozen into the solid). Therefore, the GFA is always thought to be associated with the competing crystals [38,39]. In order to make sure the relationship among the glass formation, the Zr content and the competing crystals in the
current study, extensive XRD and TEM investigations were performed. Fig. 5 compares XRD patterns of as-cast rods with 2.5 mm in diameter in the Ni61.5 Nb38.5−a Zra alloys. It indicates phase transformation dependence of Zr concentration upon solidification at the similar condition. For the alloys with the 1 at.% and 3 at.% Zr, the position of the crystalline diffraction peaks remains almost identical, but their difference only exists in the intensity. It suggests that the structures of the precipitating crystals in 1 at.% and 3 at.% Zr samples are the same. By carefully matched with the data in the Power diffraction files, the crystals were indexed as the hexagonal NiNb and orthorhombic Ni3 Nb phases, which were also confirmed in Fig. 6a and b. But with increasing the Zr content from 1 at.% to 3 at.%, the size of the crystals decreases dramatically, from −400 nm to −20 nm. The grain refinement, which causes the broadening of the peaks in XRD patterns of the 3 at.% Zr alloy as shown in Fig. 5, reveals that the Zr addition can retard the precipitation of the crystals and be conducive to the glass formation. When the
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energy E, thereby, elevating G* . As a result, it would postpone the process of the nucleation to some extent. Minor Zr below 5 at% effectively retards the nucleation and growth of the NiNb and Ni3 Nb phases and makes the grain decrease to −20 nm (Fig. 6a and b); when the Zr content is 5 at.%, the crystals are completely suppressed (Figs. 5 and 6c); but excessive Zr addition above 7 at.% prompts the separation of a new phase (Figs. 5 and 6d) so as to deteriorate the GFA [8,40]. 3.4. Mechanical property
Fig. 5. XRD patterns of the as-cast Ni61.5 Nb38.5−a Zra (a = 1, 3, . . ., 9 at.%) rods with a diameter of 2.5 mm.
Zr content rises to 5 at.%, the competing crystals were completely suppressed and the amorphous phase were formed as indicated by the unique broad diffuse halo in XRD patterns of the 5 at.% Zr alloy in Fig. 5 and the homogeneous contrast of TEM image in Fig. 6c. When the Zr content reaches 7 at.%, a new weak peak (marked by the red dash line) appears at 2 = 38.9◦ in the XRD patterns, denoting that the new primary crystalline phase is produced so as to decrease the GFA. Fig. 6d shows that some orthorhombic Ni10 (Nb, Zr)7 crystals were formed in the glassy matrix and the size is about 50 nm. When the Zr content is further added to 9 at.%, the peak at 2 = 38.9◦ in the XRD patterns in Fig. 5 largely enhanced, indicating that the sample crystallized a lot. It is consistent with Fig. 6e. The size of the grain of the primarily precipitated orthorhombic Ni10 (Nb, Zr)7 phase reaches −500 nm. Meanwhile, some unidentified phases could be formed, as marked by the capital “A” in Fig. 6e. When combining the present results with our previous work that the competing crystals against the glass formation are NiNb and Ni3 Nb phases in the binary Ni61.5 Nb38.5 alloy, it is known that introducing minor Zr below 3 at.% in the alloys of Ni61.5 Nb38.5 cannot change the type of the competing crystals but retards nucleation and growth process, which would improve the GFA, and when the Zr content increases up to 7 at.%, the new primary competing crystalline phase is produced so as to deteriorate the GFA. As a result, the alloy of the 5 at.% Zr exhibit the best GFA of the Ni61.5 Nb38.5−a Zra alloys, as shown in Figs. 1 and 3. The trend of the phase transformation with the Zr content also agrees well with the crystallization shown in Fig. 4. From a view of solidification, the crystals form through nucleation and growth of the nuclei. It would suppress the formation of the crystals to increases thermodynamically Gibbs free energy barrier G* for the nucleation. G* for a critical spherical nucleus is expressed by [40,41]: G∗ ≈
16 3 3(Gv + E)
2
In order to evaluate the mechanical performance of the studied Ni–Nb–Zr BMGs, the quasi-static compression tests were carried out. Five samples with 2 mm in diameter and 4 mm in length were measured at a strain rate of 2 × 10−4 s−1 for each alloy. Fig. 7 shows the stress curves as a function of strain of the three alloys with the highest GFA (Figs. 2 and 3). The data of mechanical properties of the alloys are tabulated in Table 2. It is seen that the measured the samples all displayed an elastic deformation up to the yield strain of 1.9–2.3% at the yield stress of about 2.7 GPa, followed by a plastic elongation by about 2% prior to the ultimate fracture. All the alloys display considerable ultimate fracture strengths as high as approximate 3.2 GPa. For a comparison, the mechanical properties of binary Ni–Nb BMGs are also listed in Table 2 [25]. It is easy to be found out that Zr addition slightly reduces the strengths of the alloys from 3.4 GPa for Ni61.5 Nb38.5 to 3.0–3.2 GPa for the Ni–Nb–Zr alloys. It is attributed to slight reduction of Tg [42]. As illustrated in Tables 1 and 2, the maximum strength, m , is proportional to Tg . The higher Tg , the higher m . It is reasonable that the Ni–Nb–Zr glassy alloys with high Tg possess high ultimate fracture strength. Besides, to our knowledge, the Ni–Nb–Zr BMG alloys are one series of those exhibiting the highest strengths in metal–metal BMGs. SEM observations indicate that the Ni–Nb–Zr BMGs mainly fractured in a shear mode and well-developed vein patterns were formed on the fractured surfaces. Some multiple shear bands are also seen on the lateral surface of the fractured specimens. However, fractographically, 9 at.% or more Zr makes the Ni–Nb–Zr BMGs transit from the ductile to the brittle [29]. 3.5. Corrosion resistance Corrosion property of the Ni–Nb–Zr bulk metallic glasses in 1 M HCl aqueous solution was investigated. No weight loss was detected for Ni–Nb–Zr BMGs after immersion in aqueous solution open to air for 1 week, indicating that the corrosion rate is very low. For a further understanding of the corrosion behaviors of Ni–Nb–Zr BMGs and studying the influence of the Zr content on the corrosion property, electrochemical measurements were performed. Fig. 8 shows the representative curves of the cathodic and anodic potentiodynamic polarization of the BMG Ni61.5 Nb38.5−a Zra alloys in 1 M HCl aqueous solution at 300 K. In evaluating the corrosion property of the materials, the most important parameters are passive region and passive current density. A wide passive region with low passive current density corresponds to the better corro-
.
Here, is the interfacial energy and Gv is Gibbs free energy difference between the crystal and liquid; E is the strain energy induced by atomic mismatch. Gv is usually negative when the melt is undercooled. In contrast, E is positive and increases with the increase of the supercooling. Apparently, when Zr is introduced into Ni–Nb alloy, Zr atoms would locate the positions which should belong to Nb or Ni atoms. Due to the large size difference between Zr and Ni or Nb atoms, it increases distinctly the atomic-level strain
Table 2 Mechanical parameters of the glassy (a) Ni61.5 Nb38.5 , (b) Ni61.5 Nb33.5 Zr5 , (c) Ni62 Nb32.3 Zr5.7 and (d) Ni63 Nb31.45 Zr5.55 at.% samples under an unaxial compressive loading at a strain rate of 2 × 10−4 s−1 Alloy
y (MPa)
εy (%)
m (MPa)
εf (%)
E (GPa)
a b c d
3000 2730 2750 2700
1.7 2.1 2.2 2.1
3450 3000 3080 3170
3.7 4.1 3.7 3.5
170 130 128 127
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Fig. 6. Bright-field TEM images of the as-cast 2.5 mm rods: (a) Ni61.5 Nb37.5 Zr1 , (b) Ni61.5 Nb35.5 Zr3 , (c) Ni61.5 Nb33.5 Zr5 , (d) Ni61.5 Nb31.5 Zr7 and (e) Ni61.5 Nb29.5 Zr9 , indicating the dependence of the microstructures on the Zr content for the Ni–Nb–Zr alloys.
sion resistance. In Fig. 8, the similar polarization behaviors were observed among the Ni61.5 Nb38.5−a Zra BMG alloys except some differences in the magnitude of the passive region and passive current density. They were spontaneously passivated with extremely wide passive region and relatively low passive current density. For a = 1, the Ni61.5 Nb38.5−a Zra glassy alloy has the passive region of approximately 1.7 V (vs. SCE), which is from −0.18 V (vs. SCE) to 1.5 V (vs. SCE), and passive current density of about 1 A m−2 . It will undergo
locally rapid dissolution when the potential exceeds 1.63 V (vs. SCE). With further increasing the Zr content, the passive region abruptly reduces to about 1.5 V (vs. SCE) and keeps stable while the passive current density declines by one or two magnitude order, which is about 0.1 A m−2 for a = 3, 5 alloys, 0.05 A m−2 for a = 7 alloy, and 0.1 A m−2 for a = 9 alloy. Accordingly, the addition of Zr is only slightly reduce the passive region but obviously decreases the passive current density, implying that the addition of the appropriate
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Fig. 7. Nominal compressive stress-strain curves of the as-cast samples with 2 mm in diameter and 4 mm in length at a strain rate of 2 × 10−4 s−1 , a, b and c for Ni63 Nb31.45 Zr5.55 , Ni62 Nb32.3 Zr5.7 and Ni61.5 Nb33.5 Zr5 , respectively.
amount of Zr is conducive to the corrosion resistance of the investigated alloys. In the potential range higher than 1.5 V (vs. SCE), the current density of the Ni61.5 Nb38.5−a Zra with a = 3–9 increases rapidly with the potential, which may be attributed to the evolution of O2 and/or Cl2 . The results indicate that the Ni–Nb–Zr BMG alloys have high corrosion resistance in the aggressive acid solution. High corrosion resistance has been regarded as one of the superior merits of metallic glasses since the discovery of amorphous Fe–Cr–P–C alloy with the extremely high anti-corrosion property [1–3,43]. Herein, to clarifying the origin of the high anticorrosion property of the current Ni–Nb–Zr BMG alloys, the surface
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Fig. 8. Potentiodynamic polarization curves of the BMG Ni61.5 Nb38.5−a Zra (a = 1,3, . . ., 9) in at.% alloys measured at a potential sweep rate of 0.333 mV/s in 1 M HCl aqueous solution open to air at 300 K.
films formed in air and in 1 M HCl solution were characterized by XPS. XPS spectra of the Ni–Nb–Zr BMG alloys consisted of the peaks of alloy elements in addition to those of oxygen and carbon. The weak Cl 2p peak was also observed on the XPS spectra of the specimens potentiodynamically polarized till 1 V (vs. SCE) in 1 M HCl solution. The C 1s peaks resulted from the unavoidable contaminant carbon on the top surface of the specimens. The O 1s spectra, shown in Fig. 9d, is comprised of the peaks arisen from the oxygen in metal–O–metal bond, metal–OH and/or bound water. The peaks
Fig. 9. XPS spectrum of the Ni61.5 Nb33.5 Zr5 BMG alloys after potentiodynamically polarized till 1 V (vs. SCE) in 1 M HCl aqueous solution open to air at 300 K: (a) Ni 2p, (b) Nb 3d, (c) Zr 3d and (d) O 1s.
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ously enhances the corrosion resistance of the Ni–Nb–Zr BMG alloys. 4. Conclusions The systematic investigations of the GFA, thermal, mechanical and corrosion properties of the Ni–Nb–Zr BMGs lead us to draw some following conclusions: (1) Bulk metallic glasses were successfully prepared in wide composition range of Ni–Nb–Zr system. There exists a wide BMG forming region, in at.%, 60 < Ni < 64, 28 < Nb < 38 and 0 < Zr < 9. The maximum diameter of the as-cast glassy rods reaches 3 mm by using copper mould injection casting method. The Zr addition improving the GFA is attributed to suppress the nucleation of the hexagonal NiNb and orthorhombic Ni3 Nb phases. (2) Ni–Nb–Zr BMGs exhibit high thermal stability with high Tg of 880–900 K and high Tx of 915–932 K. The variation of Tg is related to the crystallization behavior for the Ni61.5 Nb38.5−a Zra alloys. (3) These Ni-based BMGs exhibit good mechanical properties along with high compressive fracture strength, 3–3.2 GPa and some compressive plastic deformation of about 2%. (4) Ni–Nb–Zr BMGs possess high corrosion resistance and were spontaneously passivated with extremely wide passive region above 1.5 V (vs. SCE) and relatively low passive current density in the 1 M HCl aqueous solution open to air at 300 K, especially for Ni61.5Nb31.5Zr7, whose passive current density is on the magnitude of 10−2 A m−2 . High corrosion resistance is due to the formation of the Nb and Zr enriched surface films, which prevent the alloys from further corrosion. The Zr addition enhances the corrosion resistance by its partially taking place of metallic Ni in the surface films. Fig. 10. Cationic contents in the surface films for the Ni–Nb–Zr BMG alloy exposed to air and those potentiodynamically polarized till 1 V (vs. SCE) in 1 M HCl aqueous solution open to air at 300 K: (a) Ni61.5 Nb37.5 Zr1 and (b) Ni61.5 Nb31.5 Zr7 .
of Ni 2p, Nb 3d and Zr 3d, shown in Fig. 9a–c, respectively, correspond to their oxidized states in the surface film and their metallic states in the underlying alloy surface [31–33]. Fig. 10 shows cationic contents in the surface films for the Ni–Nb-Zr BMG alloys exposed to air and those potentiodynamically polarized till 1 V (vs. SCE) in 1 M HCl aqueous solution open to air at 300 K. The Nb and Zr were enriched in the surface films when the specimens of the investigated alloys were exposed to air. When polarized in the 1 M HCl solution, the contents of Nb and Zr in the surface film further increased. The formation of Nband Zr-enriching surface films would be responsible for the high corrosion resistance of the Ni–Nb–Zr BMG alloys, like Nb, Zr, Tienriching surface films leading to the high corrosion resistance of the Ni–Nb–Ti–Zr–Co–(Cu) glassy alloys [32,33]. In further analyzing the effect of the addition of Zr, it was found out that the Ni content was identical, about 37 at.% (shown in Fig. 10), while Zr would partially substitute for Nb in the surface films formed in air with increasing the amount of Zr. In contrast, the surface films formed in 1 M HCl solution exhibited completely different behavior with increasing the content of Zr. The Nb content maintained equal while the Ni content dramatically decreased by Zr partial substitution for Ni. Thus, with increasing the content of Zr, the decline of the content of Ni of metallic state (Fig. 9a) in the surface films is thought to contribute into the decease (shown in Fig. 8) by two magnitude order in the passive current density in the potentiodynamic polarization measurements, which obvi-
Thus, the development of the Ni–Nb–Zr BMGs with high GFA, high thermal stability, good mechanical properties and high corrosion resistance would help expand the application of the BMGs as structural materials. Acknowledgements The authors gratefully acknowledge S.J. Zheng and G.M. Cheng for the assistance of TEM experiments, and the financial support from the Ministry of Science and Technology of China (Grant Nos. 2006CB605201 and 2005DFA50860), the National Natural Science Foundation of China (Grant No. 50731005). References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15]
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Synthesis and properties of bulk metallic glasses in the ternary Ni–Nb–Zr alloy system Z.W. Zhu a,b , H.F. Zhang a,∗ , B.Z. Ding a , Z.Q. Hu a a Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang 110016, China b Graduate School of the Chinese Academy of Sciences, Beijing 100039, China
a r t i c l e
i n f o
Article history: Received 6 November 2007 Received in revised form 11 March 2008 Accepted 8 April 2008 Keywords: Bulk metallic glass Ni-based alloy Thermal property Mechanical property Corrosion resistance
a b s t r a c t Bulk metallic glasses (BMGs) with high thermal stability, good mechanical properties and high corrosion resistance were synthesized in the Ni–Nb–Zr system. A large bulk glass-forming region with 60 < Ni < 64, 28 < Nb < 38 and 0 < Zr < 9 (in at.%) was found. The critical size for the glass formation is 3 mm. These investigated Ni-based BMGs process high glass transition temperature of about 880–900 K and high onset crystallization temperature of 915–932 K as well as high compressive fracture strength of approximate 3.0–3.2 GPa along with some compressive plasticity of about 2%. Electrochemical measurements indicate they also exhibit high corrosion resistance, i.e., large passive region above 1.5 V (vs. saturated calomel reference electrode, SCE). The influence of the Zr content on the glass-forming ability (GFA) and corrosion behaviors was carefully studied, indicating that some Zr addition improves the GFA and corrosion resistance. © 2008 Elsevier B.V. All rights reserved.
1. Introduction Bulk metallic glasses (BMGs, typically referred to a minimum casting dimension larger than 1 mm) have been greatly concerned in the past few decades because they are of particular scientific and engineering interests [1–3]. Some progress in both glass-forming ability (GFA) and mechanical properties has been made, for example, the amorphous samples with the critical size over 10 mm were successfully prepared in Mg [4], Zr [5], Fe [6], Ti [7], Cu [8], Pd [9], etc., based alloys, and some Cu [10,11], Zr [12,13], Ti [7,14] based BMGs samples display very large compressive plastic strain in compression tests. It is exciting even though some problems are still puzzled. Meanwhile, due to various desirable properties, including high yield strength, hardness and elastic strain limit in addition to reasonably high fracture toughness, fatigue resistance and corrosion resistance, etc., BMGs have been tried to be made into some items such as sporting goods, surgical instruments, and strong, thin cases for electronic devices such as mobile phone and U-disc [3]. To satisfy the requirements of commercial applications, it is urgent to improve the known BMGs’ plasticity or to develop new BMGs with higher GFA and better mechanical properties, especially based on common metals, such as Al, Cu, Fe, Ni, etc.
∗ Corresponding author. Fax: +86 2423971783. E-mail address: [email protected] (H.F. Zhang). 0921-5093/$ – see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2008.04.021
In the case of Ni-based alloys, bulk metallic glasses were prepared in the complex alloy systems, such as Ni–Nb–Cr–Mo–P–B [15], Ni–Ti–Zr–(Si,Sn) [16], Ni–Nb–Ti–Zr–Co–Cu [17], Ni–Nb–Ti–Zr–Si–Sn [18,19], Ni–Nb–Sn [20,21], Ni–Cu–Ti–Zr–Al [22], Ni–Ta–Sn [23], etc. However, with compared to that of Cu-, Zr-, Ti- and Fe-, etc., based BMGs [4–9], the GFA of Ni-based amorphous alloys is a challenging subject. Because so far the maximum dimension of Ni-based BMGs’ samples is only 5 mm [22,24], the development of new Ni-based glass former with higher GFA is imperative. In the period of the conventional amorphous alloys, Ni–Nb and Ni–Zr systems were famous for their GFA. Very recently, BMG samples up to 2-mm thick were fabricated in binary Ni–Nb alloy system [25,26]. Some reported works also indicate that Ni-based Nb-bearing BMGs possess better mechanical properties than other Ni-based ones [16,27]. According to Miracle’s efficient cluster packing model [28], Ni–Nb–Zr system with good atomic size distribution (the Goldschmidt atomic radius of Zr is 0.160 nm, which is larger than 0.146 nm and 0.128 nm for Nb and Ni, respectively) could have high GFA [29,30]. Additionally, some believe that some valve metals, such as Nb and Zr, etc., enrich in the surface film to prevent the materials from corrosion [31–34]. As a result, the ternary Ni–Nb–Zr system might be a good candidate to develop new Ni-based BMGs with better combination among GFA, good mechanical properties and high corrosion resistance. Nevertheless, systematic investigation of BMGs in ternary Ni–Nb–Zr system is hardly reported [29,34,35].
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In this paper, we reported that bulk glasses can be formed in quite wide composition range of ternary Ni–Nb–Zr system. The 3mm diameter glassy samples were successfully prepared by copper mold injection casting method. They displayed good mechanical properties, high thermal stability and high anti-corrosion property. The effect of the Zr concentration on the phase transformation upon casting and corrosion behavior is also carefully discussed. 2. Experimental Master alloy ingots were prepared by arc melting a mixture of ultrasonically cleansed Ni, Nb and Zr with a purity of above 99.9% on a water-cooled copper hearth under Ti-gettered high purity argon atmosphere. The chemical homogeneity was obtained by repeated melting at least four times. The ingots were then remelted under high vacuum in a quartz tube by using induction heating coil and injected through a nozzle with 0.5–1 mm in diameter into the copper mould with a cavity of 2–4 mm diameter. The as-cast samples were characterized with X-ray diffraction (XRD, Philips PW1050, Cu K␣), transmission electron microscopy (TEM, JEOL 2010, 200 kV) and differential scanning calorimetry (DSC; Netzsch DSC 404C). The specimens used for XRD measurement were cut from the middle part of the as-cast rods. Thin slices from the 2.5-mm diameter as-cast rods were used for preparing the TEM samples, which were ground and mechanically dimpled with a GATAN precision dimple grinder as well as polished using argon ion milling as the final thinning process using a GATAN precision ion polishing system (PIPS). DSC measurements were performed in a flowing argon atmosphere at a heating rate of 0.33 K/s. Mechanical properties were measured with the samples of 2mm diameter and 4-mm length on a servo-hydraulic materials testing system (MTS 810). To perform compression tests under a constant strain rate of 2 × 10−4 s−1 , a MTS strain gauge was used. Fracture surface was examined by scanning electron microscopy (SEM, Hitachi S3400N). Corrosion behavior of Ni-based BMGs in 1 M HCl aqueous solution open to air was studied by electrochemical measurements on an Advanced Electrochemical System (Princeton Applied Research PARSTAT 2273) at 300 K. Prior to the corrosion tests, the specimens were mechanically polished in cyclohexane with silicon carbide paper up to No. 2000, degreased in acetone, washed in the distilled water, and dried in air. Electrochemical measurements were conducted in a three-electrode cell a platinum counter electrode and a saturated calomel reference electrode (SCE). Potentiodynamic polarization curves were measured at a potential sweep rate of 0.333 mV/s after immersing the samples for several minutes, when the open-circuit potential became almost steady. The surface of the samples exposed to air after mechanical polishing and conducted by potentiodynamic polarization measurements up to 1 V (vs. SCE) in 1 M HCl solution was examined by X-ray photoelectron spectroscopy (XPS) using a photoelectron spectrometer with Al K␣ radiation (h = 1486.6 eV). From these spectra, the composition of the passive film and the underlying alloy surface was quantitatively determined. 3. Results and discussion 3.1. Glass-forming ability The Ni–Nb–Zr alloy system shows good GFA. Fig. 1 illustrates XRD patterns of the as-cast Ni–Nb–Zr alloys rods with a diameter of 2 mm. In Fig. 1a, when a increase from 1 to 9, the 2-mm diameter as-cast samples of Ni61.5 Nb38.5−a Zra alloys display only a series of diffuse maxima around 2 = 42◦ , while for a = 11 the 2-mm diame-
Fig. 1. XRD patterns of the as-cast rods with a diameter of 2 mm for (a) Ni61.5 Nb38.5−a Zra (a = 1, 3, . . ., 11 at.%) [29] and (b) Ni100−b (Nb0.85 Zr0.15 )b (b = 35, 36, . . ., 40 at.%).
ter as-cast rods exhibit apparent crystalline Bragg peaks. It indicates that when Zr content is below 9 at.%, the glassy samples with 2 mm in diameter can be synthesized in the alloys Ni61.5 Nb38.5−a Zra [29]. Likewise, it is also found out that the glass can be also formed in a quite large composition range as the Ni concentration is changed. It is shown in Fig. 1b. When Ni content increases from 60 at.% to 65 at.%, the patterns of the 61–64 at.% Ni-bearing samples consist of only a broad peak without any observable crystalline diffraction peaks, indicating that the 2-mm diameter glassy sample can be made. Outside of this range, crystallization occurs to the samples with 2 mm in diameter. Identification of the crystalline phases will be discussed in detail in Section 3.3. Through tens of alloys experiments, it is discovered that there exists a wide BMG forming region of 60 < Ni < 64, 28 < Nb < 38 and 0 < Zr < 9, in at.%, in ternary Ni–Nb–Zr alloy system, as shown in Fig. 2. In the larger green ellipse region, the samples with at least 2 mm in diameter can be manufactured. It is necessary to point out that the 3-mm diameter as-cast rods are capable to be produced in the smaller purple ellipse region. XRD patterns of the 3-mm diameter samples are shown in Fig. 3, displaying the typical characteristics of those of the amorphous phase. 3.2. Thermal property Characterization of the thermal property of the investigated Ni–Nb–Zr glassy alloys, especially determination of the onset of glass transition temperature (Tg ) and the onset of crystallization
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Fig. 2. Sketch of the glass formation and variation of Tg and Tx in the Ni–Nb–Zr system. The solid green ellipse corresponds to the region where BMGs with at least 2 mm in diameter can be formed, the solid purple ellipse corresponding to 3-mm diameter BMGs.
Fig. 3. XRD patterns of the as-cast rods with a diameter of 3 mm in the three alloys.
temperature (Tx ) as well as liquidus temperature (Tl ) was conducted by DSC measurements. Fig. 4 shows the high temperature DSC profiles recorded at a heating rate of 0.33 K/s. And Table 1 lists the concerning thermal data of the typical alloys. The samples used for DSC measurements were cut from the middle section of the 2-mm diameter as-cast rods. When DSC traces and the thermal properties are considered, the alloys with the Zr content of
223
1–9 at.% or Ni content of 61–64 at.% show the distinct glass transition, while for 60 at.% Ni and 65 at.% Ni alloys, no endothermic event can be observed on the traces, implying no glass transition occurs, as shown in Fig. 4a and b. Tg , defined as the onset of the endothermic event, is relatively high and above 880 K. Compared with the known Ni-based BMGs [15–22], it can be inferred that these BMGs have high thermal stability. It is easily found out that for a series of alloys with the composition of Ni61.5 Nb38.5−a Zra , Tg decreases abruptly by about 9 K when the concentration of Zr increases from 3 at.% to 5 at.% (Fig. 4a and Table 1). When Zr content is below 3 at.% or above 5 at.%, Tg nearly remains equal, around 893 K or 882 K, respectively. The similar phenomena on glass transition does not occur to Ni100−b (Nb0.85 Zr0.15 )b alloys. Fig. 4b and Table 1 show that Tg falls continuously as Ni content decreases. Additionally, it is observed from Fig. 4a and b that there exist significant differences in crystallization behaviors. The alloys transform from the beginning three-stage crystallization to double-stage one as the Zr content or Ni content rises. Whether the Zr content increases or Ni content falls, Tx always declines. The variation of Tg and Tx is roughly drawn in Fig. 3. Further, it is found that the variation of Tg is related to the crystallization behavior for the Ni61.5 Nb38.5−a Zra alloys, since the abrupt decrease in Tg occurs at the transition from the three-stage crystallization to double-stage one. As shown in Figs. 2 and 3, the alloy with the 5 at.% Zr exhibited the best GFA of the Ni61.5 Nb38.5−a Zra alloys. As a result, the GFA is inferred to have a close relationship with the crystallization process, which will be discussed in detail in Section 3.3. It is thought that the atomic arrangement configuration is attributed to Tg or Tx dependence of Zr content. As known, the different interaction exists among Ni, Nb and Zr atoms, indicated by different mixing enthalpy values among them, i.e., −49 kJ/mol for Ni–Zr, −30 kJ/mol for Ni–Nb, and 4 kJ/mol for Nb–Zr [36]. Due to the different interaction, it leads to atomic reconfiguration to introduce Zr atoms into the Ni–Nb alloy. Extended X-ray absorption fine structure experiments [35] reveal that the bonds like Ni–Ni and Ni–Zr around Ni atoms and Nb–Ni and Nb–Nb around Nb atoms are chemically preferred to be formed as Zr is added into Ni–Nb alloys. Difference in atomic configuration contributes to the different behaviors, including different Tg , Tx , etc., during reheating process. It also affects the subsequent crystallization as mentioned above. Fig. 4c and d exemplify the melting behaviors of the Ni–Nb–Zr alloys. The liquidus temperature, Tl , and the melting temperature, Tm , decrease with increasing the Zr amount or reducing the Ni amount. The extent of the decline of Tl is faster than that of Tm . Similarly, Trg [37] deduced from the thermal parameters, are proposed to be correlated well with the GFA. They are also given in Table 1. Trg exhibits high values, but does not either possess a good correlation with the GFA in the Ni–Nb–Zr system. The supercooled liquid region T = Tx − Tg , which reflects the thermal stability of the
Table 1 Thermal properties of the as-cast samples with a diameter of 2 mm, except that the diameter of samples in Ni61.5 Nb38.5 alloy is 1.5 mm, deduced from the high temperature DSC measurement at a heating rate of 0.33 K/s Alloy (at.%)
Tg (K)
Tx (K)
Tp (K)
Tm (K)
Tl (K)
Tx (=Tx − Tg , K)
Trg (=Tg /Tl )
Ni61.5 Nb38.5−a Zra
a=0 a=1 a=3 a=5 a=7 a=9
894 893 892 883 882 882
932 926 921 918 913 913
– 935 928 926 923 918
1455 1441 1426 1414 1390 1384
1519 1510 1456 1444 1439 1420
38 33 29 35 31 31
0.589 0.591 0.613 0.612 0.613 0.620
Ni100−b (Nb0.85 Zr0.15 )b
b = 36 b = 37 b = 38 b = 39
899 890 886 884
935 933 924 915
– – – –
1408 1412 1416 1420
1448 1450 1487 1513
36 43 38 31
0.621 0.616 0.599 0.584
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Fig. 4. DSC scans corresponding to glass transitions and crystallizations, melting behaviors of the as-cast rods with a diameter of 2 mm, (a) and (c) for Ni61.5 Nb38.5−a Zra (a = 1, 3, . . ., 9 at.%) [29], (b) and (d) for Ni100−b (Nb0.85 Zr0.15 )b (b = 35, 36, . . ., 40 at.%).
supercooled liquid towards crystallization, varies slightly with the composition (Table 1) and ranges from 30 K to 40 K. It is not directly related to the GFA (shown in Fig. 3) in the current work although T was suggested to characterize the GFA [2]. 3.3. Phase transformation dependence of the Zr content upon solidification Upon solidification, the glass is formed by competing against the nucleation and growth of the crystals in the undercooled melt. The glass can be fabricated under the condition that the nucleation and growth of the primary competing crystals are completely suppressed while the melt is cooled through the temperature interval from Tl (below which crystallization is thermodynamically possible) to Tg (below which the melt is frozen into the solid). Therefore, the GFA is always thought to be associated with the competing crystals [38,39]. In order to make sure the relationship among the glass formation, the Zr content and the competing crystals in the
current study, extensive XRD and TEM investigations were performed. Fig. 5 compares XRD patterns of as-cast rods with 2.5 mm in diameter in the Ni61.5 Nb38.5−a Zra alloys. It indicates phase transformation dependence of Zr concentration upon solidification at the similar condition. For the alloys with the 1 at.% and 3 at.% Zr, the position of the crystalline diffraction peaks remains almost identical, but their difference only exists in the intensity. It suggests that the structures of the precipitating crystals in 1 at.% and 3 at.% Zr samples are the same. By carefully matched with the data in the Power diffraction files, the crystals were indexed as the hexagonal NiNb and orthorhombic Ni3 Nb phases, which were also confirmed in Fig. 6a and b. But with increasing the Zr content from 1 at.% to 3 at.%, the size of the crystals decreases dramatically, from −400 nm to −20 nm. The grain refinement, which causes the broadening of the peaks in XRD patterns of the 3 at.% Zr alloy as shown in Fig. 5, reveals that the Zr addition can retard the precipitation of the crystals and be conducive to the glass formation. When the
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energy E, thereby, elevating G* . As a result, it would postpone the process of the nucleation to some extent. Minor Zr below 5 at% effectively retards the nucleation and growth of the NiNb and Ni3 Nb phases and makes the grain decrease to −20 nm (Fig. 6a and b); when the Zr content is 5 at.%, the crystals are completely suppressed (Figs. 5 and 6c); but excessive Zr addition above 7 at.% prompts the separation of a new phase (Figs. 5 and 6d) so as to deteriorate the GFA [8,40]. 3.4. Mechanical property
Fig. 5. XRD patterns of the as-cast Ni61.5 Nb38.5−a Zra (a = 1, 3, . . ., 9 at.%) rods with a diameter of 2.5 mm.
Zr content rises to 5 at.%, the competing crystals were completely suppressed and the amorphous phase were formed as indicated by the unique broad diffuse halo in XRD patterns of the 5 at.% Zr alloy in Fig. 5 and the homogeneous contrast of TEM image in Fig. 6c. When the Zr content reaches 7 at.%, a new weak peak (marked by the red dash line) appears at 2 = 38.9◦ in the XRD patterns, denoting that the new primary crystalline phase is produced so as to decrease the GFA. Fig. 6d shows that some orthorhombic Ni10 (Nb, Zr)7 crystals were formed in the glassy matrix and the size is about 50 nm. When the Zr content is further added to 9 at.%, the peak at 2 = 38.9◦ in the XRD patterns in Fig. 5 largely enhanced, indicating that the sample crystallized a lot. It is consistent with Fig. 6e. The size of the grain of the primarily precipitated orthorhombic Ni10 (Nb, Zr)7 phase reaches −500 nm. Meanwhile, some unidentified phases could be formed, as marked by the capital “A” in Fig. 6e. When combining the present results with our previous work that the competing crystals against the glass formation are NiNb and Ni3 Nb phases in the binary Ni61.5 Nb38.5 alloy, it is known that introducing minor Zr below 3 at.% in the alloys of Ni61.5 Nb38.5 cannot change the type of the competing crystals but retards nucleation and growth process, which would improve the GFA, and when the Zr content increases up to 7 at.%, the new primary competing crystalline phase is produced so as to deteriorate the GFA. As a result, the alloy of the 5 at.% Zr exhibit the best GFA of the Ni61.5 Nb38.5−a Zra alloys, as shown in Figs. 1 and 3. The trend of the phase transformation with the Zr content also agrees well with the crystallization shown in Fig. 4. From a view of solidification, the crystals form through nucleation and growth of the nuclei. It would suppress the formation of the crystals to increases thermodynamically Gibbs free energy barrier G* for the nucleation. G* for a critical spherical nucleus is expressed by [40,41]: G∗ ≈
16 3 3(Gv + E)
2
In order to evaluate the mechanical performance of the studied Ni–Nb–Zr BMGs, the quasi-static compression tests were carried out. Five samples with 2 mm in diameter and 4 mm in length were measured at a strain rate of 2 × 10−4 s−1 for each alloy. Fig. 7 shows the stress curves as a function of strain of the three alloys with the highest GFA (Figs. 2 and 3). The data of mechanical properties of the alloys are tabulated in Table 2. It is seen that the measured the samples all displayed an elastic deformation up to the yield strain of 1.9–2.3% at the yield stress of about 2.7 GPa, followed by a plastic elongation by about 2% prior to the ultimate fracture. All the alloys display considerable ultimate fracture strengths as high as approximate 3.2 GPa. For a comparison, the mechanical properties of binary Ni–Nb BMGs are also listed in Table 2 [25]. It is easy to be found out that Zr addition slightly reduces the strengths of the alloys from 3.4 GPa for Ni61.5 Nb38.5 to 3.0–3.2 GPa for the Ni–Nb–Zr alloys. It is attributed to slight reduction of Tg [42]. As illustrated in Tables 1 and 2, the maximum strength, m , is proportional to Tg . The higher Tg , the higher m . It is reasonable that the Ni–Nb–Zr glassy alloys with high Tg possess high ultimate fracture strength. Besides, to our knowledge, the Ni–Nb–Zr BMG alloys are one series of those exhibiting the highest strengths in metal–metal BMGs. SEM observations indicate that the Ni–Nb–Zr BMGs mainly fractured in a shear mode and well-developed vein patterns were formed on the fractured surfaces. Some multiple shear bands are also seen on the lateral surface of the fractured specimens. However, fractographically, 9 at.% or more Zr makes the Ni–Nb–Zr BMGs transit from the ductile to the brittle [29]. 3.5. Corrosion resistance Corrosion property of the Ni–Nb–Zr bulk metallic glasses in 1 M HCl aqueous solution was investigated. No weight loss was detected for Ni–Nb–Zr BMGs after immersion in aqueous solution open to air for 1 week, indicating that the corrosion rate is very low. For a further understanding of the corrosion behaviors of Ni–Nb–Zr BMGs and studying the influence of the Zr content on the corrosion property, electrochemical measurements were performed. Fig. 8 shows the representative curves of the cathodic and anodic potentiodynamic polarization of the BMG Ni61.5 Nb38.5−a Zra alloys in 1 M HCl aqueous solution at 300 K. In evaluating the corrosion property of the materials, the most important parameters are passive region and passive current density. A wide passive region with low passive current density corresponds to the better corro-
.
Here, is the interfacial energy and Gv is Gibbs free energy difference between the crystal and liquid; E is the strain energy induced by atomic mismatch. Gv is usually negative when the melt is undercooled. In contrast, E is positive and increases with the increase of the supercooling. Apparently, when Zr is introduced into Ni–Nb alloy, Zr atoms would locate the positions which should belong to Nb or Ni atoms. Due to the large size difference between Zr and Ni or Nb atoms, it increases distinctly the atomic-level strain
Table 2 Mechanical parameters of the glassy (a) Ni61.5 Nb38.5 , (b) Ni61.5 Nb33.5 Zr5 , (c) Ni62 Nb32.3 Zr5.7 and (d) Ni63 Nb31.45 Zr5.55 at.% samples under an unaxial compressive loading at a strain rate of 2 × 10−4 s−1 Alloy
y (MPa)
εy (%)
m (MPa)
εf (%)
E (GPa)
a b c d
3000 2730 2750 2700
1.7 2.1 2.2 2.1
3450 3000 3080 3170
3.7 4.1 3.7 3.5
170 130 128 127
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Fig. 6. Bright-field TEM images of the as-cast 2.5 mm rods: (a) Ni61.5 Nb37.5 Zr1 , (b) Ni61.5 Nb35.5 Zr3 , (c) Ni61.5 Nb33.5 Zr5 , (d) Ni61.5 Nb31.5 Zr7 and (e) Ni61.5 Nb29.5 Zr9 , indicating the dependence of the microstructures on the Zr content for the Ni–Nb–Zr alloys.
sion resistance. In Fig. 8, the similar polarization behaviors were observed among the Ni61.5 Nb38.5−a Zra BMG alloys except some differences in the magnitude of the passive region and passive current density. They were spontaneously passivated with extremely wide passive region and relatively low passive current density. For a = 1, the Ni61.5 Nb38.5−a Zra glassy alloy has the passive region of approximately 1.7 V (vs. SCE), which is from −0.18 V (vs. SCE) to 1.5 V (vs. SCE), and passive current density of about 1 A m−2 . It will undergo
locally rapid dissolution when the potential exceeds 1.63 V (vs. SCE). With further increasing the Zr content, the passive region abruptly reduces to about 1.5 V (vs. SCE) and keeps stable while the passive current density declines by one or two magnitude order, which is about 0.1 A m−2 for a = 3, 5 alloys, 0.05 A m−2 for a = 7 alloy, and 0.1 A m−2 for a = 9 alloy. Accordingly, the addition of Zr is only slightly reduce the passive region but obviously decreases the passive current density, implying that the addition of the appropriate
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Fig. 7. Nominal compressive stress-strain curves of the as-cast samples with 2 mm in diameter and 4 mm in length at a strain rate of 2 × 10−4 s−1 , a, b and c for Ni63 Nb31.45 Zr5.55 , Ni62 Nb32.3 Zr5.7 and Ni61.5 Nb33.5 Zr5 , respectively.
amount of Zr is conducive to the corrosion resistance of the investigated alloys. In the potential range higher than 1.5 V (vs. SCE), the current density of the Ni61.5 Nb38.5−a Zra with a = 3–9 increases rapidly with the potential, which may be attributed to the evolution of O2 and/or Cl2 . The results indicate that the Ni–Nb–Zr BMG alloys have high corrosion resistance in the aggressive acid solution. High corrosion resistance has been regarded as one of the superior merits of metallic glasses since the discovery of amorphous Fe–Cr–P–C alloy with the extremely high anti-corrosion property [1–3,43]. Herein, to clarifying the origin of the high anticorrosion property of the current Ni–Nb–Zr BMG alloys, the surface
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Fig. 8. Potentiodynamic polarization curves of the BMG Ni61.5 Nb38.5−a Zra (a = 1,3, . . ., 9) in at.% alloys measured at a potential sweep rate of 0.333 mV/s in 1 M HCl aqueous solution open to air at 300 K.
films formed in air and in 1 M HCl solution were characterized by XPS. XPS spectra of the Ni–Nb–Zr BMG alloys consisted of the peaks of alloy elements in addition to those of oxygen and carbon. The weak Cl 2p peak was also observed on the XPS spectra of the specimens potentiodynamically polarized till 1 V (vs. SCE) in 1 M HCl solution. The C 1s peaks resulted from the unavoidable contaminant carbon on the top surface of the specimens. The O 1s spectra, shown in Fig. 9d, is comprised of the peaks arisen from the oxygen in metal–O–metal bond, metal–OH and/or bound water. The peaks
Fig. 9. XPS spectrum of the Ni61.5 Nb33.5 Zr5 BMG alloys after potentiodynamically polarized till 1 V (vs. SCE) in 1 M HCl aqueous solution open to air at 300 K: (a) Ni 2p, (b) Nb 3d, (c) Zr 3d and (d) O 1s.
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ously enhances the corrosion resistance of the Ni–Nb–Zr BMG alloys. 4. Conclusions The systematic investigations of the GFA, thermal, mechanical and corrosion properties of the Ni–Nb–Zr BMGs lead us to draw some following conclusions: (1) Bulk metallic glasses were successfully prepared in wide composition range of Ni–Nb–Zr system. There exists a wide BMG forming region, in at.%, 60 < Ni < 64, 28 < Nb < 38 and 0 < Zr < 9. The maximum diameter of the as-cast glassy rods reaches 3 mm by using copper mould injection casting method. The Zr addition improving the GFA is attributed to suppress the nucleation of the hexagonal NiNb and orthorhombic Ni3 Nb phases. (2) Ni–Nb–Zr BMGs exhibit high thermal stability with high Tg of 880–900 K and high Tx of 915–932 K. The variation of Tg is related to the crystallization behavior for the Ni61.5 Nb38.5−a Zra alloys. (3) These Ni-based BMGs exhibit good mechanical properties along with high compressive fracture strength, 3–3.2 GPa and some compressive plastic deformation of about 2%. (4) Ni–Nb–Zr BMGs possess high corrosion resistance and were spontaneously passivated with extremely wide passive region above 1.5 V (vs. SCE) and relatively low passive current density in the 1 M HCl aqueous solution open to air at 300 K, especially for Ni61.5Nb31.5Zr7, whose passive current density is on the magnitude of 10−2 A m−2 . High corrosion resistance is due to the formation of the Nb and Zr enriched surface films, which prevent the alloys from further corrosion. The Zr addition enhances the corrosion resistance by its partially taking place of metallic Ni in the surface films. Fig. 10. Cationic contents in the surface films for the Ni–Nb–Zr BMG alloy exposed to air and those potentiodynamically polarized till 1 V (vs. SCE) in 1 M HCl aqueous solution open to air at 300 K: (a) Ni61.5 Nb37.5 Zr1 and (b) Ni61.5 Nb31.5 Zr7 .
of Ni 2p, Nb 3d and Zr 3d, shown in Fig. 9a–c, respectively, correspond to their oxidized states in the surface film and their metallic states in the underlying alloy surface [31–33]. Fig. 10 shows cationic contents in the surface films for the Ni–Nb-Zr BMG alloys exposed to air and those potentiodynamically polarized till 1 V (vs. SCE) in 1 M HCl aqueous solution open to air at 300 K. The Nb and Zr were enriched in the surface films when the specimens of the investigated alloys were exposed to air. When polarized in the 1 M HCl solution, the contents of Nb and Zr in the surface film further increased. The formation of Nband Zr-enriching surface films would be responsible for the high corrosion resistance of the Ni–Nb–Zr BMG alloys, like Nb, Zr, Tienriching surface films leading to the high corrosion resistance of the Ni–Nb–Ti–Zr–Co–(Cu) glassy alloys [32,33]. In further analyzing the effect of the addition of Zr, it was found out that the Ni content was identical, about 37 at.% (shown in Fig. 10), while Zr would partially substitute for Nb in the surface films formed in air with increasing the amount of Zr. In contrast, the surface films formed in 1 M HCl solution exhibited completely different behavior with increasing the content of Zr. The Nb content maintained equal while the Ni content dramatically decreased by Zr partial substitution for Ni. Thus, with increasing the content of Zr, the decline of the content of Ni of metallic state (Fig. 9a) in the surface films is thought to contribute into the decease (shown in Fig. 8) by two magnitude order in the passive current density in the potentiodynamic polarization measurements, which obvi-
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