Hot Consolidation of Partially Amorphous Cu-Ti Based Alloy: a Comparison Between Hot Extrusion and Hot Compaction by Sintering

Abstract

Consolidation of amorphous powders, which take advantage of the supercooled liquid region, is an alternative way to overcome the size limitation in marginal metallic glasses. Cu₃₆Ti₃₄Zr₂₂Ni₈based amorphous powders were obtained during high energy ball-milling. The analyses revealed that amorphous structures of powders and ribbons are quite different and this led to a different thermal behavior being the powders more thermally stable than the ribbons. Extrusion was initially proposed as a consolidation process; however, the decrease in viscosity in remainder amorphous matrix was not sufficient for that process, but certainly sufficient for sintering the sample during hot consolidation. An amorphous/nanocrystalline microstructure known for enhancing the mechanical properties of their fully amorphous counterparts was obtained. Evaluation of mechanical properties by microhardness revealed the relatively high hardness of HV 768. From these results, consolidation by sintering seems a promising route to produce bulk metallic glasses nanocomposites.

nanocomposites; consolidation; amorphous alloys

1 Introduction

Cu-Ti based bulk metallic glasses (BMG) alloys have attracted interest due to their superior mechanical properties and relatively lower costs^1-4. However, the lack of ductility has limited the applications. To overcome this problem, composites of embedded nanocrystals in the Cu-Ti glassy matrix are being developed.^{5, 6}. Furthermore, normal casting routes present the issue of size limitation, in which the largest obtained thickness was around 5mm³. The route of solid state amorphization followed by powder consolidation appears to be an attractive route to produce sized BMGs in alloy systems with marginal glass forming ability, as is the case of Cu-Ti-based alloys. Processes of amorphization by mechanical alloying have been described in Cu-Ti-Zr-Ni systems^7-9. Consolidation of amorphous powder can be achieved by ultra-high pressure (UHP), by taking advantage of the decrease in viscosity of the supercooled-liquid region¹⁰, sintering-spark plasma (SPS)^11-13, equal channel angular extrusion (ECAE)¹⁴ and extrusion¹⁵. According to preliminary tests¹⁶ , the Cu₃₆Ti₃₄Zr₂₂Ni₈ alloy is a strong candidate to form metallic glass. It has been shown¹⁶ that this alloy presents amorphous phase formation when processed by melt spinning showing supercooled-liquid region (ΔT_x=T_x–T_g), ΔT_X = 67K¹⁶. The Assuming that Cu and Ni are similar, Cu₃₆Ti₃₄Zr₂₂Ni₈ composition is located near to deep eutectic in a Cu-Ti-Zr ternary phase diagram, which decreases the probability of intermetallic phases precipitation¹⁷. Therefore, in the present paper, the Cu₃₆Ti₃₄Zr₂₂Ni₈ alloy was chosen to produce an amorphous structure by ball milling followed by hot compaction and sintering.

2 Experimental

Cu₃₆Ti₃₄Zr₂₂Ni₈ alloy ingots were prepared by arc-melting of high purity constituent elements (i.e., Cu (99,99%), Ti (99,99%), and Zr (99,50%), Ni (99,99%)) in a high-purity argon gas atmosphere and using Ti-gettering. All the ingots were re-melted at least 4 times to obtain a homogeneous master alloy. The ingots were crushed in a hammer-forging die before milling. After crushing, the material was loaded together with hardened-steel balls in a hardened steel vial to perform milling with a ball-to-powder ratio of 30:1. The ball milling process was performed in a Zoz Simoloyer with water cooling temperature control, at 900 rpm, under purified argon gas atmosphere during 8 and 16 hours. The obtained powders were cold compacted at room temperature under 1GPa of uniaxial pressure. Finally, samples with a diameter of 4.5 mm were produced by hot-extrusion, which was done inside the supercooled liquid region at 736 K, with ram speed of 1 mm/min and an extrusion ratio of 3:1. The structural characterizations of powder and extruded samples were performed by X-ray diffraction (DRX) in a Siemens D5005 diffractometer with CuKα radiation, and transmission electron microscopy (TEM) in a Tecnai F20 Philips. The glass transition (T_g) and crystallization (T_x) temperatures were determined by differential scanning calorimetry in a Netzsch DSC 200F3 Maia at a rate of 0.67 K/s. The mechanical properties were analyzed by using a Omnimet MHT Buehler – Wilson Hardness Tukon 1102 system with Vickers indenter loaded by 500 g. The system is automatic and applies the load during a time following ASTM E384¹⁸, and ISO 6507-1¹⁹ standards.

3 Results and Discussion

Figure 1 shows the comparison between the XRD patterns of ribbon¹⁶, 8h and 16h ball-milled powders, and crystalline powders. This comparison shows that ball-milled powders did not present well-defined crystalline peaks, and that both peak-broadening and peak-positions are not matched to the ones for the ribbon. This can be an indication that the mechanism of amorphization is not the same. It has been reported that during ball-milling of pre-alloyed powders, the amorphization is thought to occur by the accumulation of structural defects such as vacancies, dislocations, grain boundaries, and anti-phase boundaries²⁰. In comparison with amorphous alloys obtained by rapid solidification, the amorphization through reactions in the solid state has, in principle, the advantage to lead to techniques for producing bulky amorphous materials²¹. Moreover, due to the processing characteristics, it is reasonable to imagine that the physical properties, such as Tg and Tx of an amorphous alloy produced by the solid state reaction during milling, may be different from those of the amorphous alloy prepared by rapid cooling²¹.

Figure 1
X-ray diffraction patterns of melt-spun ribbon, ball-milled powders (8h and 16h) and crystalline powders of Cu₃₆Ti₃₄Zr₂₂Ni alloy.

In accordance with what was expected from the phase diagram, the unique phase precipitated in crystalline powder is Cu₂TiZr and the solid solution of Cu with Ni shifted the peaks slightly to right. If one compares the ball-milled powder with the crystalline ones, the small peaks appearing at about 30, 35 and 63 degrees (the more intense) can be assigned as being the zirconium oxide, produced during milling. Due to the presence of higher intense peaks in the crystalline powders, ZrO does not appear in the spectrum and this is an indication that ZrO comprises less than 5% of the whole matrix.

Figure 2 presents the DSC curves of ball-milled powder in comparison with the ribbon, before correction of internal relaxation. The DSC curves confirm the formation of a totally different amorphous phase in the powders, which presented a higher glass transition temperature (T_g) and onset of the primary crystallization (T_x) than the ribbons. Moreover the T_g and T_x for ribbons are respectively about 100 K and 72 K lower than the ones for ball-milled powders. This will reflect in a larger supercooled liquid region for ribbons, ΔT_x = 67K, when compared with the ball-milled powders, i.e., ΔT_x = 39K and ΔT_x = 36K, respectively for 16 and 8 hours. Even with lower ΔTx, the highest Tx observed for the powders means highest thermal stability²², which may be good for final mechanical applications. However, this will also create difficulties for further powder hot-consolidation due to the reduction of the window processing.

Figure 2
DSC curves for melt-spun ribbon and for ball-milled powders (8h and 16h) of Cu₃₆Ti₃₄Zr₂₂Ni₈alloy.

Figure 3 shows the bright field TEM images (BF-TEM) and selected area electron diffraction patterns (SAEDP) for both 8h and 16h ball-milled powders. From these images, it is not possible to observe the presence of nanocrystals, because they are BF images whose contrast can be influenced by differences in thickness. However, the SAEDP of powders milled till 8h is showing the presence of an almost continuous ring formed by small white spots inside a diffused halo, indicating the presence of very small crystals inside of an amorphous structure. The SAEDPs for powders milled for 16 h is showing a broader diffraction halo, when compared with 8 h of milling, as well as continuous outer rings formed by very small spots. The indexation of these outer halos indicated the presence of ZrO, which increased with the milling time. The increased broadening and the absence of in-halo spots indicate that the sample is totally or, at least for the BF condition, "electron-diffraction" amorphous.

Figure 3
Bright field image and SAEDP of Cu₃₆Ti₃₄Ni₈Zr₂₂ ball-milled powders. (a) 8h and (b) 16h.

These nanocrystals can be better observed in Figure 4 which presents images of high resolution transmission electron microscopy (HRTEM) of the same regions presented in Figure 3. Figure 4a clearly shows cross-grating of Moiré lattice fringes, which are highlighted by red dashed circles and zoomed in the lower right inset, also confirmed by the Fast Fourier Transform (FFT) in the upper right inset, which presents a very high amount of spots. Even considering that they were barely observed, these lattice fringes indicate the precipitation of nanocrystals within the amorphous matrix²³. The sizes of these nanocrystals were estimated to be about 20 nm. Figure 4b does not show clear evidence of nanocrystallization and the FFT presented in the inset clearly shows a pure amorphous state, which indicates that the effect of additional milling time resulted in a greater amorphous amount.

Figure 4
HRTEM images of Cu₃₆Ti₃₄Ni₈Zr₂₂ ball-milled powders (a) 8h and (b) 16h.

Moreover, the reinforced composite BMG/nanocrystals have been proven effective in improving their mechanical properties⁵. Even the extrusion process has been done in the supercooled liquid region, just a small part of the sample could be fully extruded (Figure 5a). This can be explained by effective nanocrystallization during the passage in the conical region (partially extruded), where the material is subjected to an intense shearing and locally generates heat^{24, 25}. This can induce localized nanocrystallization which increases the viscosity. In this context, the pressure needed to continue the process became too high for the pressure machine and it was impossible to continue the process. Further, the extruded part presented brittle behavior, and it was not possible to cut the sample for further preparation. Nevertheless, the previous part of the sample was intensely hot compacted by simple sinterization (Figure 5b) and the amorphous matrix was preserved. This can be observed in Figure 6a where the XRD pattern shows that the amorphous matrix is maintained inside the hot consolidated sample, but there was also a noticeable nanocrystallization indicated by an evolution of the intensity of crystalline peaks. The maintenance of amorphous matrix is also confirmed by the DSC curve presented in Figure 6b which shows crystallization and an increase of both T_g and T_x temperatures, resulting in a ΔT_x of 40K.

Figure 5
(a) Hot extruded sample; (b) hot consolidated of 16h ball-milled Cu₃₆Ti₃₄Zr₂₂Ni powders.

Figure 6
(a) X-ray diffraction pattern and; (b) DSC curve for 16h ball-milled hot consolidated of Cu₃₆Ti₃₄Zr₂₂Ni alloy.

Figure 7 shows images of BF-TEM, SAEDP in the inset of Figure 7a, and HRTEM of 16h ball-milled hot consolidated sample (Figure 7b). As was the case in Figure 3, from Figure 7a it is not possible to observe the presence of nanocrystals, but, differently from SAEDPs of Figure 3, the SAEDP of Figure 7 shows the presence of a continuous ring formed by small white spots inside a diffused halo. This means that crystals with very small sizes are embedded in the amorphous matrix and that these sizes are even smaller than the ones in Figure 3. These nanocrystals can be better observed in Figure 7b which presents HRTEM images of the same region presented in Figure 7a. Figure 7b clearly shows cross-grating of Moiré lattice fringes in the upper-right inset of the Figure, and this is clear evidence of precipitation of nanocrystals within the amorphous matrix. The sizes of these nanocrystals were estimated to be about 2 nm.

Figure 7
(a) Bright field image and SAED pattern; (b) HRTEM image of 16h ball-milled hot consolidated Cu₃₆Ti₃₄Ni₈Zr₂₂powders.

The hardness test results show that the hot consolidated sample reaches HV 768. If one can assume that the yield strength (ρ) is ρ≈HV/3,6^{14, 26} this will give ρ ≈ 2765 MPa. These values can be considered relatively high when compared with others billets obtained in a consolidation process of Cu-Ti based alloys¹³.

From these results, it is apparent that traditional powder metallurgy followed by sinterization can be successfully used to produce sized BMGs or, at least, composites of BMG matrix and nanocrystals which can produce materials with enhanced mechanical properties.

4 Summary and Conclusions

Amorphous/crystalline nanocomposite powders were produced by high energy ball-milling and consolidated during hot extrusion process. Microstructural, thermal and mechanical investigations were carried out.

The microstructural and thermal analysis revealed the different mechanism of amorphization of the powders when compared with ribbons. Powders milled for 16h presented highest thermal stability indicated by highest crystallization temperature.

The amorphous matrix with embedded nanocrystals microstructure was maintained during extrusion. The hot consolidated samples presented a hardness value of 768 HV which is relatively higher than for other Cu-Ti based billets produced by similar consolidation process.

Traditional powder metallurgy followed by sinterization is a promising route to produce sized BMGs with enhanced mechanical properties.

Acknowledgements

The authors thank the National Council for Scientific and Technological Development (CNPq - Brazil) for a doctoral grant and Foundation for Research Support of the State of São Paulo (FAPESP - Brazil) - Thematic Project for financial support.

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Publication Dates

• Publication in this collection
  May-Jun 2015

History

• Received
  10 Mar 2014
• Reviewed
  08 Apr 2015