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Enhanced glass forming ability in Zr-based bulk metallic glasses with Hf Addition
Intermetallics 24 (2012) 115e119
Contents lists available at SciVerse ScienceDirect
Intermetallics journal homepage: www.elsevier.com/locate/intermet
Enhanced glass forming ability in Zr-based bulk metallic glasses with Hf Addition Dongchun Qiao*, Atakan Peker Applied Sciences Laboratory, Institute for Shock Physics, Washington State University, 120 N Pine Street, P. O. 1495, Spokane, WA 99210, USA
a r t i c l e i n f o
a b s t r a c t
Article history: Received 23 December 2011 Received in revised form 25 January 2012 Accepted 31 January 2012 Available online 25 February 2012
We report on the partial substitution of Hf for Zr in Zr57Nb5Cu15.4Ni12.6Al10 bulk metallic glass (BMG) and the resulting enhanced glass forming ability and processability. Critical casting thickness increased from less than 14 mm for Zr57Nb5Cu15.4Ni12.6Al10 to more than 16 mm for Zr45Hf12Nb5Cu15.4Ni12.6Al10. This improvement contrasts with earlier reports of reduced glass forming ability with partial substitution of Hf for Zr. Furthermore, the new BMG alloy exhibited one of the best bend ductility reported to date for BMG samples that are more than 4 mm thick. The improved glass forming ability with Hf substitution is presented and the underlying factors for improved glass forming ability are discussed. Ó 2012 Elsevier Ltd. All rights reserved.
Keywords: B. Glasses, metallic C. Casting D: Microstructure
1. Introduction Bulk Metallic Glasses (BMGs) have received significant attention in academia and industry over the past two decades. While the mechanical properties of BMGs (e.g. very high strength and large elastic limit) are beneficial for engineering applications, the accessibility into deeply undercooled liquid regime makes BMGs valuable for the study of liquid and glass structure in academic research as well [1,2]. Because the chemical formulation of BMGs is crucial both to their engineering properties and their glass forming ability, numerous BMG formulations have been developed since the 1990’s [2e4]. Of these, Zr-based BMGs are the most well-developed and commonly used system in both fundamental research and commercial applications [5]. High processability and large critical casting thicknesses are among the leading attributes of Zr-based BMGs. Casting thicknesses of greater than 16 mm for Zr41.2Ti13.8Ni10.0Cu12.5Be22.5, 16 mm for Zr55Al10Ni5Cu30, and 12 mm for Zr57Nb5Cu15.4Ni12.6Al10 have been reported [3,5e8]. A variety of elements have been added or substituted into Zrbased BMGs to improve their glass forming ability (GFA) and engineering properties [8e10]. Hf, however, has not been a preferred element for this purpose given that a substitution of more than 5% was reported to significantly reduce glass forming ability [11e13]. We report enhanced GFA and processibility with a 12% substitution of Hf for Zr in the Zr-based BMG
* Corresponding author. Tel.: þ1 509 358 7844; fax: þ1 509 358 7700. E-mail address: [email protected] (D. Qiao). 0966-9795/$ e see front matter Ó 2012 Elsevier Ltd. All rights reserved. doi:10.1016/j.intermet.2012.01.028
Zr57Nb5Cu15.4Ni12.6Al10 (Vit-106). The critical casting thickness was increased from less than 14 mm for Zr57Nb5Cu15.4Ni12.6Al10 to greater than 16 mm for Zr45Hf12Nb5Cu15.4Ni12.6Al10. 2. Experimental procedures Master alloys of Zr57Nb5Cu15.4Ni12.6Al10, Zr45Hf12Nb5Cu15.4Ni12.6Al10, and Zr28.5Hf28.5Nb5Cu15.4Ni12.6Al10 were prepared by fusing the elemental metals Zr, Hf, Cu, Ni, Al, and Nb with purities of 99.5%e99.95%. The Zr and Hf metals were crystal bar grade. The nominal oxygen content of these metals was less than 100 ppm. Master alloy buttons were made using a laboratory arc-melter on a water-cooled copper hearth in a Ti-gettered argon atmosphere. A two-step melting procedure was used. First, Ni and Nb in each composition were fused into a binary Ni-Nb master alloy, which was then fused with the other elemental metals. Each melting step was repeated several times with sample flipping to ensure chemical homogeneity. The arc-melted samples were formed into either in circular buttons or elongated ingots. Cylindrical rod samples were prepared by remelting the above arc-melted samples in a 316 stainless steel tubular container under vacuum, followed by waterquenching. The stainless steel tubing and a nominal reaction layer were removed using lathe machining. The samples were sliced along the cross-section using an abrasive saw for optical microscopy (OM), hardness tests, X-ray Diffraction (XRD) and Differential Scanning Calorimetry (DSC). Hardness tests were performed using an automated LECO Rockwell instrument with a load of 60 kg f (Rockwell A scale). The X-ray diffraction pattern was collected using X’Pert Cu Ka radiation with a wave length of 1.54 Å. Differential
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scanning calorimetry was performed using a Netzsch 409 Differential Scanning Calorimetry (DSC) in an Al2O3 crucible using a 20 K/min heating rate in a flowing Ar atmosphere. 3. Results 3.1. Glass forming ability in master alloy buttons Pictures of as-prepared master alloy buttons for Zr57Nb5Cu15.4Ni12.6Al10, Zr45Hf12Nb5Cu15.4Ni12.6Al10, and Zr28.5Hf28.5Nb5Cu15.4Ni12.6Al10 are presented in Fig. 1 (all alloy formulations are given in atomic percent). Fig. 1(a), (b), and (c) are 30 g master alloy buttons of Zr57Nb5Cu15.4Ni12.6Al10, Zr45Hf12Nb5Cu15.4Ni12.6Al10, and Zr28.5Hf28.5Nb5Cu15.4Ni12.6Al10, respectively. The buttons were primarily cooled from the bottom, which was in contact with the copper hearth. Thus, the upper portion of a button had the lowest cooling rate. Only Zr45Hf12Nb5Cu15.4Ni12.6Al10 alloy showed a mirror-like luster and surface smoothness, indicating substantial glass formation on the upper portion of the button. The Zr57Nb5Cu15.4Ni12.6Al10 alloy exhibited a minor level of sink and surface roughness. Zr28.5Hf28.5Nb5Cu15.4Ni12.6Al10 had the most sink and surface roughness, indicating significant crystallization on its upper portion. Given these results, 40 g buttons of Zr57Nb5Cu15.4Ni12.6Al10 and Zr45Hf12Nb5Cu15.4Ni12.6Al10 were prepared to confirm the trend. The 40 g master alloy button of Zr45Hf12Nb5Cu15.4Ni12.6Al10 (Fig. 1(e)) still had a mirror-like luster and surface smoothness, whereas the 40 g master alloy button of Zr57Nb5Cu15.4Ni12.6Al10 (Fig. 1(d)) exhibited a larger level of sink and surface roughness, indicating much more crystallization than that of the 30 g button. Accordingly, our preliminary conclusion from these observations was that a modest substitution of Hf for Zr, 12% atomic in this case, enhanced GFA, whereas the much larger Hf substitution of 28.5% degraded GFA of the Zr57Nb5Cu15.4Ni12.6Al10 BMG. These observations were also based on prior observations that volume shrinkage during metallic glass formation is much less than the volume shrinkage associated with the nucleation and growth of crystalline phases [14] Optical microscopy and hardness tests were also performed on the master alloy buttons to confirm the observations from visual examination of the buttons. Cross-sectional views of the Zr57Nb5Cu15.4Ni12.6Al10 and Zr45Hf12Nb5Cu15.4Ni12.6Al10 buttons are presented in Fig. 2. The pictures were taken approximately from the mid-center of the master alloy buttons and show both the representative microstructure and indentation marks from Rockwell hardness testing. The 30 g master alloy button of Zr28.5Hf28.5Nb5Cu15.4Ni12.6Al10 was so brittle that an acceptable sample for metallographic mounting and polishing could not be readily obtained. This alloy was therefore not analyzed or tested further. Optical microscopy revealed a featureless microstructure in the 30 g master alloy buttons of both Zr57Nb5Cu15.4Ni12.6Al10 and Zr45Hf12Nb5Cu15.4Ni12.6Al10, indicating good glass formation
(Fig. 2(a) and (c)). Furthermore, the indentation marks from Rockwell hardness testing are associated with out-arching shear bands, a trademark of the high quality amorphous phase of BMG. [5]. Given the 60 kgf load in the Rockwell hardness testing and the absence of crack formation, the hardness indentations indicate formation of a relatively tough amorphous phase in both buttons. Of the 40 g alloy buttons, Zr45Hf12Nb5Cu15.4Ni12.6Al10 still exhibited a featureless microstructure and out-arching shear bands around the hardness indentation (Fig. 2(d)). The 40 g Zr57Nb5Cu15.4Ni12.6Al10 button (Fig. 2(b)), however, exhibited a rather dramatic change with the featureless microstructure replaced by a polycrystalline grainy microstructure. Furthermore, the hardness indentations caused several large cracks without any out-arching shear bands, proving that the 40 g Zr57Nb5Cu15.4Ni12.6Al10 button had either a minimal amount of amorphous phase or none at all. These observations further confirm the improved GFA of Zr45Hf12Nb5Cu15.4Ni12.6Al10 over Zr57Nb5Cu15.4Ni12.6Al10. 3.2. Cylindrical rod castings Another sample production method was used to positively confirm the enhanced glass forming ability resulting from partial substitution of Hf for Zr. Cylindrical rod samples were prepared by remelting the master alloy buttons in vacuum-sealed stainless steel tubular containers which were submerged in a controlled heat bath and then water quenched. The cylindrical rod length was at least four times its diameter to minimize the effect of fast cooling from the end. This approach yields more controlled processing conditions, including casting temperature, and a more consistent cooling rate. Samples of Zr57Nb5Cu15.4Ni12.6Al10 and Zr45Hf12Nb5Cu15.4Ni12.6Al10 were prepared using stainless steel tubing with an outer diameter of 12.7 mm. This yielded BMG samples of approximately 11 mm in diameter. Both alloys produced fully amorphous samples confirmed by both optical microscopy and X-ray diffraction. The next set of BMG samples were prepared using stainless steel tubing with an outer diameter of 15.9 mm to produce BMG samples approximately 14 mm in diameter. The 14 mm diameter cylindrical rod of Zr45Hf12Nb5Cu15.4Ni12.6Al10 yielded a fully amorphous sample, as confirmed by both optical microscopy and X-ray diffraction. However, the 14 mm diameter cylindrical rod of Zr57Nb5Cu15.4Ni12.6Al10 began to show some crystalline phases in a primarily amorphous matrix. Fig. 3(a) presents the optical microscopy and Rockwell hardness indentation of this sample imaged at the center of a cross-sectional view. The microstructure reveals a featureless phase (amorphous) combined with a clustered precipitation of crystalline phases. These crystalline phases are currently unidentified. Furthermore, the Rockwell hardness indentation caused both our-arching shear bands and small cracks, as seen in red circles in Fig. 3(a). This clearly indicates that the upper bound of glass forming ability in the Zr57Nb5Cu15.4Ni12.6Al10 alloy is reached around 14 mm in diameter. Finally, a still larger diameter
Fig. 1. Top surface of master alloy buttons. (a) 30 g Zr57Nb5Cu15.4Ni12.6Al10, (b) 30 g Zr45Hf12Nb5Cu15.4Ni12.6Al10, (c) 30 g Zr28.5Hf28.5Nb5Cu15.4Ni12.6Al10, (d) 40 g Zr57Nb5Cu15.4Ni12.6Al10 and (e) 40 g Zr45Hf12Nb5Cu15.4Ni12.6Al10.
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Fig. 2. Polished cross-section of master alloy buttons and indentation marks from Rockwell hardness testing. (a) 30 g Zr57Nb5Cu15.4Ni12.6Al10, (b) 40 g Zr57Nb5Cu15.4Ni12.6Al10, (c) 30 g Zr45Hf12Nb5Cu15.4Ni12.6Al10 and (d) 40 g Zr45Hf12Nb5Cu15.4Ni12.6Al10.
alloy, Zr45Hf12Nb5Cu15.4Ni12.6Al10, was prepared using stainless steel tubing with an outer diameter of 19 mm. This produced a BMG sample approximately 16 mm in diameter. Fig. 3(b) presents the optical microscopy and Rockwell hardness indentation of this sample imaged at the center of a cross-sectional view. The microstructure reveals a clean and featureless phase (amorphous) and out-arching shear bands around the hardness indentation, indicating that the upper bound of GFA in the Zr45Hf12Nb5Cu15.4Ni12.6Al10 alloy is not yet reached at a diameter of 16 mm. XRD patterns were collected from these last two samples at the center of a cross-sectional view of a 14 mm diameter rod of Zr57Nb5Cu15.4Ni12.6Al10 and a 16 mm diameter rod of Zr45Hf12Nb5Cu15.4Ni12.6Al10. Fig. 4(a) and (b) depicts these XRD patterns. Note that the pattern consists of only a broad diffraction peak without a detectable sharp Bragg peak for the 16 mm diameter Zr45Hf12Nb5Cu15.4Ni12.6Al10 rod, indicating the fully amorphous structure of the sample within the XRD resolution (Fig. 4(b)). On the other hand, there are some notable sharp peaks superimposed on the broad peak of the 14 mm diameter Zr57Nb5Cu15.4Ni12.6Al10 rod,
indicating some crystalline phases, as yet unidentified, formed during solidification (Fig. 4(a)). These XRD results are in good agreement with the results of the optical microscopy and Rockwell hardness tests. They show, positively, that a 12% substitution of Hf for Zr in Zr57Nb5Cu15.4Ni12.6Al10 can significantly enhance GFA. Using the same preparation and processing method, the critical casting thickness increased from less than 14 mm for Zr57Nb5Cu15.4Ni12.6Al10 to more than 16 mm for Zr45Hf12Nb5Cu15.4Ni12.6Al10. This is a notable improvement, particularly considering earlier reports that the substitution of Hf for Zr diminished GFA. [11e13] 3.3. Thermal analysis Differential Scanning Calorimetry (DSC) measurements were performed to determine characteristic thermal properties, glass transition temperature (Tg), crystallization temperature (Tx) temperature, and the melting temperature of crystalline phases (Tm). The DSC samples were cut from 11 mm diameter glassy rod of
Fig. 3. Polished cross-section of cylindrical rod castings and indentation marks from Rockwell hardness testing. (a) 14 mm diameter Zr57Nb5Cu15.4Ni12.6Al10 and (b) 16 mm diameter Zr45Hf12Nb5Cu15.4Ni12.6Al10. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
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Fig. 4. X-ray diffraction pattern of cylindrical rod castings. (a) 14 mm diameter Zr57Nb5Cu15.4Ni12.6Al10 and (b) 16 mm diameter Zr45Hf12Nb5Cu15.4Ni12.6Al10.
Fig. 5. DSC scans of BMG samples run at 20 K/min (a) 11 mm diameter Zr57Nb5Cu15.4Ni12.6Al10 and (b) 14 mm diameter Zr45Hf12Nb5Cu15.4Ni12.6Al10.
Zr57Nb5Cu15.4Ni12.6Al10 and 14 mm diameter glassy rod of Zr45Hf12Nb5Cu15.4Ni12.6Al10. Both DSC scans (Fig. 5), were run at 20 K/min, and exhibited a clear endothermic event due to glass transition, followed by exothermic events indicating transformations from the supercooled liquid regime to crystalline phases. On continuous heating, much larger endothermic events were observed, indicating the onset of melting. Given the relative sharpness and shape of the melting curves, the onset of melting temperatures gives the eutectic temperature for these glass forming alloys. Accordingly, these onset temperatures were used as melting temperatures Tm in reduced glass transition Trg calculations. Similarly, the onset temperatures were used for Tg and Tx. As noted in Fig. 5, the substitution of Hf for Zr increases both Tg, and Tm. This is expected due to the higher refractory nature of Hf and is consistent with prior reports [11,12]. Meanwhile, the reduced glass transition temperature (Trg ¼ Tg/Tm) minimally changed from 0.628 for Zr57Nb5Cu15.4Ni12.6Al10 to 0.625 for Zr45Hf12Nb5Cu15.4Ni12.6Al10. Given the various assumptions of Turnbull’s glass forming criteria and possible experimental errors, such a small difference in Trg values is not enough to predict the GFA enhancement from Zr57Nb5Cu15.4Ni12.6Al10 to Zr45Hf12Nb5Cu15.4Ni12.6Al10.
substitution. Moreover, the reports suggested that Hf substitution may potentially degrade GFA in Zr-based BMGs based on the experimentally determined “critical casting thickness” or decreasing values of Trg. For example, Figueroa [13] et al reported a dramatic drop in the critical casting thickness of HfxZr45xCu45Ag10, from 3.5 mm down to 1 mm, when x is increased from 0 to 5. Though the same study reported some GFA improvement with 1.5% Hf substitution for Zr, no explanation is provided. Certainly, Hf substitution of more than 5% resulted in an unchanged GFA at best, and sometimes led to substantial GFA degradation. Our results bring an interesting twist to the effect of various elements on the glass forming ability of Zr-based BMGs, most notably in support of the so called “confusion principle” of BMG alloy design [16]. It is generally accepted that any new element added to glass forming alloys should have a significant difference in atomic size to make the confusion principle work effectively [16]. Given the fact that Hf and Zr have the same atomic radii (155 picometers) and very similar chemical properties, the improvement in GFA we achieved with partial substitution of Hf for Zr is significant for the confusion principle. It appears that a new element added to glass forming alloys does not need to have a significant difference in atomic size to improve GFA, provided the amount of alloying addition is carefully selected. Adding more elements and increasing the complexity of alloy formulation helps to stabilize the liquid structure against crystalline phases by increasing its entropy. Thus,
3.4. 3-point bend test Various 3-point bend test samples were prepared from 11 mm diameter glassy rods of Zr45Hf12Nb5Cu15.4Ni12.6Al10. These samples showed significant bend ductility in 3-point bend testing. One of these samples, after 3-point bend testing, is shown in Fig. 6. The sample has a rectangular cross-section 5.0 mm thick and 7.6 mm wide. During this test, the sample was subjected to 8% plastic strain at its maximum deflection point without failure. A large portion of the sample, where the elastic strain limit is exceeded, exhibited global plastic deformation and bend ductility without any observable crack formation. To the best of our knowledge, this is the largest BMG sample to show permanent deformation under bending; it is also the largest plastic strain reported for BMG samples that are more than 4 mm thick [15]. More results on the 3-point bend testing and the bend ductility of Zr45Hf12Nb5Cu15.4Ni12.6Al10 BMG will be reported in a separate paper. 4. Discussion Earlier reports [11e13] on the substitution of Hf for Zr in Zrbased BMGs found no GFA enhancement above 5% partial
Fig. 6. The permanent deformation of a 3-point bend test sample of Zr45Hf12Nb5Cu15.4Ni12.6Al10 BMG with a rectangular cross-section of 5.0 mm thick and 7.6 mm wide. The sample was subjected to 8% plastic strain at its maximum deflection point.
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a significant difference in atomic size in alloying additions is not a necessary condition for the confusion principle to work effectively. However, the catch is that no new crystallization mechanism (such as the formation of new crystalline phases absent in simpler systems) becomes effective with increased alloy complexity to dominate the positive effect of the confusion principle in glass forming ability. We note several reports on enhanced GFA in BMGs with minute additions of alloying elements. For example, a fraction of a percent, or a few percent of Y [17] (or Si or Hf, as mentioned above), can improve the glass forming ability of certain BMGs. The high level of GFA enhancement from minor alloying additions is not truly compatible with the confusion principle. Although the mechanisms for such improvements are not fully understood, the positive effect of these minute alloying additions on GFA is best explained by the fluxing effect on heterogeneous nucleation. 5. Conclusions In summary, the substitution of Hf for Zr in Zr57Nb5Cu15.4Ni12.6Al10 leads to improved GFA. The critical casting size increases from less than 14 mm for Zr57Nb5Cu15.4Ni12.6Al10 to greater than 16 mm for Zr45Hf12Nb5Cu15.4Ni12.6Al10. Furthermore, the new BMG alloy exhibited one of the best bend ductility reported to date for BMG samples that are more than 4 mm thick. Acknowledgments This work was supported by the U.S. Office of Naval Research (Grant #N00014-06-1-0315). The authors thank Nate Arganbright for his technical assistance in the XRD measurements.
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References [1] Johnson WL. Bulk glass-forming metallic alloys: science and technology. MRS Bulletin (USA) 1999;24(10):42e56. [2] Inoue A. Stabilization of metallic supercooled liquid and bulk amorphous alloys. Acta Materialia 2000;48(1):279e306. [3] Peker A, Johnson WL. A highly processable metallic-glass Zr41.2Ti13.8Ni10.0Cu12.5Be22.5. Applied Physics Letters 1993;63(17):2342e4. [4] Lin XH, Johnson WL. Formation of Ti-Zr-Cu-Ni bulk metallic glasses. Journal of Applied Physics 1995;78(11):6514e9. [5] Peker A. Science, technology & applications of bulk amorphous alloys. In: International Workshop on Bulk Metallic Glasses: Science and Technology, Bangalore-560012, India, January 12e16, 2007. [6] Inoue A, Zhang T. Fabrication of bulky Zr-based glassy alloys by suction casting into copper mold. Materials Transaction JIM 1995;36(9):1184e7. [7] Lin XH, Johnson WL, Rhim WK. Effect of oxygen impurity on crystallization of an undercooled bulk glass forming Zr-Ti-Cu-Ni-Al alloy. Materials Transaction JIM 1997;38(5):473e7. [8] Lin XH. Ph.D. Thesis: California Institute of Technology; 1997. [9] Wang WH, Bian Z, Wen P, Zhang Y, Pan MX, Zhao DQ. Role of addition in formation and properties of Zr-based bulk metallic glasses. Intermetallics 2002;10(11e12):1249e57. [10] Lu ZP, Liu CT. Role of minor alloying additions in formation of bulk metallic glasses: a review. Journal of Materials Science 2004;39(12):3965e74. [11] Gu X, Xing L, Hufnagel T. Glass-forming ability and crystallization of bulk metallic glass (HfxZr1x) 52.5Cu17.9Ni14.6Al10Ti5. Journal of Non-Crystalline Solids 2002;311:77e82. [12] Gu X, Jiao T, Kecskes LJ, Woodman RH, Fan C, Ramesh KT, et al. Crystallization and mechanical behavior of (Hf, Zr)-Ti-Cu-Ni-Al metallic glasses. Journal of Non-Crystalline Solids 2003;317:112e7. [13] Figueroa IA, Zhao H, Gonzalez S, Davis HA, Todd I. Bulk glass formability for Cu-Hf-Zr-Ag and Cu-Zr-Ag-Si alloys. Journal of Non-Crystalline Solids 2008; 354:5181e3. [14] Kui HW, Greer AL, Turnbull D. Formation of bulk metallic glass by fluxing. Applied Physics Letters 1984;45:615. [15] Schroers J, Johnson WL. Ductile bulk metallic glass. Physical Review Letters 2004;93(25):255506. [16] Greer AL. Materials science-confusion design. Nature 1993;366:303e4. [17] Wang WH. Roles of minor additions in formation and properties of bulk metallic glass. Progress in Materials Science 2007;52:540e96.
Contents lists available at SciVerse ScienceDirect
Intermetallics journal homepage: www.elsevier.com/locate/intermet
Enhanced glass forming ability in Zr-based bulk metallic glasses with Hf Addition Dongchun Qiao*, Atakan Peker Applied Sciences Laboratory, Institute for Shock Physics, Washington State University, 120 N Pine Street, P. O. 1495, Spokane, WA 99210, USA
a r t i c l e i n f o
a b s t r a c t
Article history: Received 23 December 2011 Received in revised form 25 January 2012 Accepted 31 January 2012 Available online 25 February 2012
We report on the partial substitution of Hf for Zr in Zr57Nb5Cu15.4Ni12.6Al10 bulk metallic glass (BMG) and the resulting enhanced glass forming ability and processability. Critical casting thickness increased from less than 14 mm for Zr57Nb5Cu15.4Ni12.6Al10 to more than 16 mm for Zr45Hf12Nb5Cu15.4Ni12.6Al10. This improvement contrasts with earlier reports of reduced glass forming ability with partial substitution of Hf for Zr. Furthermore, the new BMG alloy exhibited one of the best bend ductility reported to date for BMG samples that are more than 4 mm thick. The improved glass forming ability with Hf substitution is presented and the underlying factors for improved glass forming ability are discussed. Ó 2012 Elsevier Ltd. All rights reserved.
Keywords: B. Glasses, metallic C. Casting D: Microstructure
1. Introduction Bulk Metallic Glasses (BMGs) have received significant attention in academia and industry over the past two decades. While the mechanical properties of BMGs (e.g. very high strength and large elastic limit) are beneficial for engineering applications, the accessibility into deeply undercooled liquid regime makes BMGs valuable for the study of liquid and glass structure in academic research as well [1,2]. Because the chemical formulation of BMGs is crucial both to their engineering properties and their glass forming ability, numerous BMG formulations have been developed since the 1990’s [2e4]. Of these, Zr-based BMGs are the most well-developed and commonly used system in both fundamental research and commercial applications [5]. High processability and large critical casting thicknesses are among the leading attributes of Zr-based BMGs. Casting thicknesses of greater than 16 mm for Zr41.2Ti13.8Ni10.0Cu12.5Be22.5, 16 mm for Zr55Al10Ni5Cu30, and 12 mm for Zr57Nb5Cu15.4Ni12.6Al10 have been reported [3,5e8]. A variety of elements have been added or substituted into Zrbased BMGs to improve their glass forming ability (GFA) and engineering properties [8e10]. Hf, however, has not been a preferred element for this purpose given that a substitution of more than 5% was reported to significantly reduce glass forming ability [11e13]. We report enhanced GFA and processibility with a 12% substitution of Hf for Zr in the Zr-based BMG
* Corresponding author. Tel.: þ1 509 358 7844; fax: þ1 509 358 7700. E-mail address: [email protected] (D. Qiao). 0966-9795/$ e see front matter Ó 2012 Elsevier Ltd. All rights reserved. doi:10.1016/j.intermet.2012.01.028
Zr57Nb5Cu15.4Ni12.6Al10 (Vit-106). The critical casting thickness was increased from less than 14 mm for Zr57Nb5Cu15.4Ni12.6Al10 to greater than 16 mm for Zr45Hf12Nb5Cu15.4Ni12.6Al10. 2. Experimental procedures Master alloys of Zr57Nb5Cu15.4Ni12.6Al10, Zr45Hf12Nb5Cu15.4Ni12.6Al10, and Zr28.5Hf28.5Nb5Cu15.4Ni12.6Al10 were prepared by fusing the elemental metals Zr, Hf, Cu, Ni, Al, and Nb with purities of 99.5%e99.95%. The Zr and Hf metals were crystal bar grade. The nominal oxygen content of these metals was less than 100 ppm. Master alloy buttons were made using a laboratory arc-melter on a water-cooled copper hearth in a Ti-gettered argon atmosphere. A two-step melting procedure was used. First, Ni and Nb in each composition were fused into a binary Ni-Nb master alloy, which was then fused with the other elemental metals. Each melting step was repeated several times with sample flipping to ensure chemical homogeneity. The arc-melted samples were formed into either in circular buttons or elongated ingots. Cylindrical rod samples were prepared by remelting the above arc-melted samples in a 316 stainless steel tubular container under vacuum, followed by waterquenching. The stainless steel tubing and a nominal reaction layer were removed using lathe machining. The samples were sliced along the cross-section using an abrasive saw for optical microscopy (OM), hardness tests, X-ray Diffraction (XRD) and Differential Scanning Calorimetry (DSC). Hardness tests were performed using an automated LECO Rockwell instrument with a load of 60 kg f (Rockwell A scale). The X-ray diffraction pattern was collected using X’Pert Cu Ka radiation with a wave length of 1.54 Å. Differential
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scanning calorimetry was performed using a Netzsch 409 Differential Scanning Calorimetry (DSC) in an Al2O3 crucible using a 20 K/min heating rate in a flowing Ar atmosphere. 3. Results 3.1. Glass forming ability in master alloy buttons Pictures of as-prepared master alloy buttons for Zr57Nb5Cu15.4Ni12.6Al10, Zr45Hf12Nb5Cu15.4Ni12.6Al10, and Zr28.5Hf28.5Nb5Cu15.4Ni12.6Al10 are presented in Fig. 1 (all alloy formulations are given in atomic percent). Fig. 1(a), (b), and (c) are 30 g master alloy buttons of Zr57Nb5Cu15.4Ni12.6Al10, Zr45Hf12Nb5Cu15.4Ni12.6Al10, and Zr28.5Hf28.5Nb5Cu15.4Ni12.6Al10, respectively. The buttons were primarily cooled from the bottom, which was in contact with the copper hearth. Thus, the upper portion of a button had the lowest cooling rate. Only Zr45Hf12Nb5Cu15.4Ni12.6Al10 alloy showed a mirror-like luster and surface smoothness, indicating substantial glass formation on the upper portion of the button. The Zr57Nb5Cu15.4Ni12.6Al10 alloy exhibited a minor level of sink and surface roughness. Zr28.5Hf28.5Nb5Cu15.4Ni12.6Al10 had the most sink and surface roughness, indicating significant crystallization on its upper portion. Given these results, 40 g buttons of Zr57Nb5Cu15.4Ni12.6Al10 and Zr45Hf12Nb5Cu15.4Ni12.6Al10 were prepared to confirm the trend. The 40 g master alloy button of Zr45Hf12Nb5Cu15.4Ni12.6Al10 (Fig. 1(e)) still had a mirror-like luster and surface smoothness, whereas the 40 g master alloy button of Zr57Nb5Cu15.4Ni12.6Al10 (Fig. 1(d)) exhibited a larger level of sink and surface roughness, indicating much more crystallization than that of the 30 g button. Accordingly, our preliminary conclusion from these observations was that a modest substitution of Hf for Zr, 12% atomic in this case, enhanced GFA, whereas the much larger Hf substitution of 28.5% degraded GFA of the Zr57Nb5Cu15.4Ni12.6Al10 BMG. These observations were also based on prior observations that volume shrinkage during metallic glass formation is much less than the volume shrinkage associated with the nucleation and growth of crystalline phases [14] Optical microscopy and hardness tests were also performed on the master alloy buttons to confirm the observations from visual examination of the buttons. Cross-sectional views of the Zr57Nb5Cu15.4Ni12.6Al10 and Zr45Hf12Nb5Cu15.4Ni12.6Al10 buttons are presented in Fig. 2. The pictures were taken approximately from the mid-center of the master alloy buttons and show both the representative microstructure and indentation marks from Rockwell hardness testing. The 30 g master alloy button of Zr28.5Hf28.5Nb5Cu15.4Ni12.6Al10 was so brittle that an acceptable sample for metallographic mounting and polishing could not be readily obtained. This alloy was therefore not analyzed or tested further. Optical microscopy revealed a featureless microstructure in the 30 g master alloy buttons of both Zr57Nb5Cu15.4Ni12.6Al10 and Zr45Hf12Nb5Cu15.4Ni12.6Al10, indicating good glass formation
(Fig. 2(a) and (c)). Furthermore, the indentation marks from Rockwell hardness testing are associated with out-arching shear bands, a trademark of the high quality amorphous phase of BMG. [5]. Given the 60 kgf load in the Rockwell hardness testing and the absence of crack formation, the hardness indentations indicate formation of a relatively tough amorphous phase in both buttons. Of the 40 g alloy buttons, Zr45Hf12Nb5Cu15.4Ni12.6Al10 still exhibited a featureless microstructure and out-arching shear bands around the hardness indentation (Fig. 2(d)). The 40 g Zr57Nb5Cu15.4Ni12.6Al10 button (Fig. 2(b)), however, exhibited a rather dramatic change with the featureless microstructure replaced by a polycrystalline grainy microstructure. Furthermore, the hardness indentations caused several large cracks without any out-arching shear bands, proving that the 40 g Zr57Nb5Cu15.4Ni12.6Al10 button had either a minimal amount of amorphous phase or none at all. These observations further confirm the improved GFA of Zr45Hf12Nb5Cu15.4Ni12.6Al10 over Zr57Nb5Cu15.4Ni12.6Al10. 3.2. Cylindrical rod castings Another sample production method was used to positively confirm the enhanced glass forming ability resulting from partial substitution of Hf for Zr. Cylindrical rod samples were prepared by remelting the master alloy buttons in vacuum-sealed stainless steel tubular containers which were submerged in a controlled heat bath and then water quenched. The cylindrical rod length was at least four times its diameter to minimize the effect of fast cooling from the end. This approach yields more controlled processing conditions, including casting temperature, and a more consistent cooling rate. Samples of Zr57Nb5Cu15.4Ni12.6Al10 and Zr45Hf12Nb5Cu15.4Ni12.6Al10 were prepared using stainless steel tubing with an outer diameter of 12.7 mm. This yielded BMG samples of approximately 11 mm in diameter. Both alloys produced fully amorphous samples confirmed by both optical microscopy and X-ray diffraction. The next set of BMG samples were prepared using stainless steel tubing with an outer diameter of 15.9 mm to produce BMG samples approximately 14 mm in diameter. The 14 mm diameter cylindrical rod of Zr45Hf12Nb5Cu15.4Ni12.6Al10 yielded a fully amorphous sample, as confirmed by both optical microscopy and X-ray diffraction. However, the 14 mm diameter cylindrical rod of Zr57Nb5Cu15.4Ni12.6Al10 began to show some crystalline phases in a primarily amorphous matrix. Fig. 3(a) presents the optical microscopy and Rockwell hardness indentation of this sample imaged at the center of a cross-sectional view. The microstructure reveals a featureless phase (amorphous) combined with a clustered precipitation of crystalline phases. These crystalline phases are currently unidentified. Furthermore, the Rockwell hardness indentation caused both our-arching shear bands and small cracks, as seen in red circles in Fig. 3(a). This clearly indicates that the upper bound of glass forming ability in the Zr57Nb5Cu15.4Ni12.6Al10 alloy is reached around 14 mm in diameter. Finally, a still larger diameter
Fig. 1. Top surface of master alloy buttons. (a) 30 g Zr57Nb5Cu15.4Ni12.6Al10, (b) 30 g Zr45Hf12Nb5Cu15.4Ni12.6Al10, (c) 30 g Zr28.5Hf28.5Nb5Cu15.4Ni12.6Al10, (d) 40 g Zr57Nb5Cu15.4Ni12.6Al10 and (e) 40 g Zr45Hf12Nb5Cu15.4Ni12.6Al10.
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Fig. 2. Polished cross-section of master alloy buttons and indentation marks from Rockwell hardness testing. (a) 30 g Zr57Nb5Cu15.4Ni12.6Al10, (b) 40 g Zr57Nb5Cu15.4Ni12.6Al10, (c) 30 g Zr45Hf12Nb5Cu15.4Ni12.6Al10 and (d) 40 g Zr45Hf12Nb5Cu15.4Ni12.6Al10.
alloy, Zr45Hf12Nb5Cu15.4Ni12.6Al10, was prepared using stainless steel tubing with an outer diameter of 19 mm. This produced a BMG sample approximately 16 mm in diameter. Fig. 3(b) presents the optical microscopy and Rockwell hardness indentation of this sample imaged at the center of a cross-sectional view. The microstructure reveals a clean and featureless phase (amorphous) and out-arching shear bands around the hardness indentation, indicating that the upper bound of GFA in the Zr45Hf12Nb5Cu15.4Ni12.6Al10 alloy is not yet reached at a diameter of 16 mm. XRD patterns were collected from these last two samples at the center of a cross-sectional view of a 14 mm diameter rod of Zr57Nb5Cu15.4Ni12.6Al10 and a 16 mm diameter rod of Zr45Hf12Nb5Cu15.4Ni12.6Al10. Fig. 4(a) and (b) depicts these XRD patterns. Note that the pattern consists of only a broad diffraction peak without a detectable sharp Bragg peak for the 16 mm diameter Zr45Hf12Nb5Cu15.4Ni12.6Al10 rod, indicating the fully amorphous structure of the sample within the XRD resolution (Fig. 4(b)). On the other hand, there are some notable sharp peaks superimposed on the broad peak of the 14 mm diameter Zr57Nb5Cu15.4Ni12.6Al10 rod,
indicating some crystalline phases, as yet unidentified, formed during solidification (Fig. 4(a)). These XRD results are in good agreement with the results of the optical microscopy and Rockwell hardness tests. They show, positively, that a 12% substitution of Hf for Zr in Zr57Nb5Cu15.4Ni12.6Al10 can significantly enhance GFA. Using the same preparation and processing method, the critical casting thickness increased from less than 14 mm for Zr57Nb5Cu15.4Ni12.6Al10 to more than 16 mm for Zr45Hf12Nb5Cu15.4Ni12.6Al10. This is a notable improvement, particularly considering earlier reports that the substitution of Hf for Zr diminished GFA. [11e13] 3.3. Thermal analysis Differential Scanning Calorimetry (DSC) measurements were performed to determine characteristic thermal properties, glass transition temperature (Tg), crystallization temperature (Tx) temperature, and the melting temperature of crystalline phases (Tm). The DSC samples were cut from 11 mm diameter glassy rod of
Fig. 3. Polished cross-section of cylindrical rod castings and indentation marks from Rockwell hardness testing. (a) 14 mm diameter Zr57Nb5Cu15.4Ni12.6Al10 and (b) 16 mm diameter Zr45Hf12Nb5Cu15.4Ni12.6Al10. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
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Fig. 4. X-ray diffraction pattern of cylindrical rod castings. (a) 14 mm diameter Zr57Nb5Cu15.4Ni12.6Al10 and (b) 16 mm diameter Zr45Hf12Nb5Cu15.4Ni12.6Al10.
Fig. 5. DSC scans of BMG samples run at 20 K/min (a) 11 mm diameter Zr57Nb5Cu15.4Ni12.6Al10 and (b) 14 mm diameter Zr45Hf12Nb5Cu15.4Ni12.6Al10.
Zr57Nb5Cu15.4Ni12.6Al10 and 14 mm diameter glassy rod of Zr45Hf12Nb5Cu15.4Ni12.6Al10. Both DSC scans (Fig. 5), were run at 20 K/min, and exhibited a clear endothermic event due to glass transition, followed by exothermic events indicating transformations from the supercooled liquid regime to crystalline phases. On continuous heating, much larger endothermic events were observed, indicating the onset of melting. Given the relative sharpness and shape of the melting curves, the onset of melting temperatures gives the eutectic temperature for these glass forming alloys. Accordingly, these onset temperatures were used as melting temperatures Tm in reduced glass transition Trg calculations. Similarly, the onset temperatures were used for Tg and Tx. As noted in Fig. 5, the substitution of Hf for Zr increases both Tg, and Tm. This is expected due to the higher refractory nature of Hf and is consistent with prior reports [11,12]. Meanwhile, the reduced glass transition temperature (Trg ¼ Tg/Tm) minimally changed from 0.628 for Zr57Nb5Cu15.4Ni12.6Al10 to 0.625 for Zr45Hf12Nb5Cu15.4Ni12.6Al10. Given the various assumptions of Turnbull’s glass forming criteria and possible experimental errors, such a small difference in Trg values is not enough to predict the GFA enhancement from Zr57Nb5Cu15.4Ni12.6Al10 to Zr45Hf12Nb5Cu15.4Ni12.6Al10.
substitution. Moreover, the reports suggested that Hf substitution may potentially degrade GFA in Zr-based BMGs based on the experimentally determined “critical casting thickness” or decreasing values of Trg. For example, Figueroa [13] et al reported a dramatic drop in the critical casting thickness of HfxZr45xCu45Ag10, from 3.5 mm down to 1 mm, when x is increased from 0 to 5. Though the same study reported some GFA improvement with 1.5% Hf substitution for Zr, no explanation is provided. Certainly, Hf substitution of more than 5% resulted in an unchanged GFA at best, and sometimes led to substantial GFA degradation. Our results bring an interesting twist to the effect of various elements on the glass forming ability of Zr-based BMGs, most notably in support of the so called “confusion principle” of BMG alloy design [16]. It is generally accepted that any new element added to glass forming alloys should have a significant difference in atomic size to make the confusion principle work effectively [16]. Given the fact that Hf and Zr have the same atomic radii (155 picometers) and very similar chemical properties, the improvement in GFA we achieved with partial substitution of Hf for Zr is significant for the confusion principle. It appears that a new element added to glass forming alloys does not need to have a significant difference in atomic size to improve GFA, provided the amount of alloying addition is carefully selected. Adding more elements and increasing the complexity of alloy formulation helps to stabilize the liquid structure against crystalline phases by increasing its entropy. Thus,
3.4. 3-point bend test Various 3-point bend test samples were prepared from 11 mm diameter glassy rods of Zr45Hf12Nb5Cu15.4Ni12.6Al10. These samples showed significant bend ductility in 3-point bend testing. One of these samples, after 3-point bend testing, is shown in Fig. 6. The sample has a rectangular cross-section 5.0 mm thick and 7.6 mm wide. During this test, the sample was subjected to 8% plastic strain at its maximum deflection point without failure. A large portion of the sample, where the elastic strain limit is exceeded, exhibited global plastic deformation and bend ductility without any observable crack formation. To the best of our knowledge, this is the largest BMG sample to show permanent deformation under bending; it is also the largest plastic strain reported for BMG samples that are more than 4 mm thick [15]. More results on the 3-point bend testing and the bend ductility of Zr45Hf12Nb5Cu15.4Ni12.6Al10 BMG will be reported in a separate paper. 4. Discussion Earlier reports [11e13] on the substitution of Hf for Zr in Zrbased BMGs found no GFA enhancement above 5% partial
Fig. 6. The permanent deformation of a 3-point bend test sample of Zr45Hf12Nb5Cu15.4Ni12.6Al10 BMG with a rectangular cross-section of 5.0 mm thick and 7.6 mm wide. The sample was subjected to 8% plastic strain at its maximum deflection point.
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a significant difference in atomic size in alloying additions is not a necessary condition for the confusion principle to work effectively. However, the catch is that no new crystallization mechanism (such as the formation of new crystalline phases absent in simpler systems) becomes effective with increased alloy complexity to dominate the positive effect of the confusion principle in glass forming ability. We note several reports on enhanced GFA in BMGs with minute additions of alloying elements. For example, a fraction of a percent, or a few percent of Y [17] (or Si or Hf, as mentioned above), can improve the glass forming ability of certain BMGs. The high level of GFA enhancement from minor alloying additions is not truly compatible with the confusion principle. Although the mechanisms for such improvements are not fully understood, the positive effect of these minute alloying additions on GFA is best explained by the fluxing effect on heterogeneous nucleation. 5. Conclusions In summary, the substitution of Hf for Zr in Zr57Nb5Cu15.4Ni12.6Al10 leads to improved GFA. The critical casting size increases from less than 14 mm for Zr57Nb5Cu15.4Ni12.6Al10 to greater than 16 mm for Zr45Hf12Nb5Cu15.4Ni12.6Al10. Furthermore, the new BMG alloy exhibited one of the best bend ductility reported to date for BMG samples that are more than 4 mm thick. Acknowledgments This work was supported by the U.S. Office of Naval Research (Grant #N00014-06-1-0315). The authors thank Nate Arganbright for his technical assistance in the XRD measurements.
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