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Crystalline to amorphous transformation in Zr–Cu–Al alloys induced by high pressure torsion
Intermetallics 37 (2013) 52e58
Contents lists available at SciVerse ScienceDirect
Intermetallics journal homepage: www.elsevier.com/locate/intermet
Crystalline to amorphous transformation in ZreCueAl alloys induced by high pressure torsion F.Q. Meng a, b, *, K. Tsuchiya a, b, Y. Yokoyama c a
Graduate School of Pure and Applied Sciences, University of Tsukuba, Tsukuba, Ibaraki 305-8577, Japan Structural Materials Unit, National Institute for Materials Science, Tsukuba, Ibaraki 305-0047, Japan c Institute for Materials Research, Tohoku University, Sendai, Miyagi 980-8577, Japan b
a r t i c l e i n f o
a b s t r a c t
Article history: Received 27 November 2012 Received in revised form 15 January 2013 Accepted 22 January 2013 Available online 26 February 2013
ZreCueAl alloys with various Al contents were processed by high pressure torsion (HPT). It was found that HPT can induce a remarkable crystalline to amorphous transformation in these alloys, with an almost fully amorphous structure obtained after 50 revolutions. Microstructural investigations indicated that the transformation from crystalline to amorphous structure is strongly dependent on both the accumulative torsional strain and chemical composition. Higher accumulative shear strain and Al content led to more pronounced amorphization. The amorphization occurred preferentially in the martensite phases with twin structure in Zr50Cu40Al10. Differential scanning calorimetric (DSC) analysis revealed a similar glass transition temperature in the amorphous samples produced by HPT to those prepared by rapid quenching. Ó 2013 Elsevier Ltd. All rights reserved.
Keywords: A. Multiphase intermetallics B. Glasses, metallic D. Microstructure D. Martensitic structure
1. Introduction Amorphous alloys are characterized by the absence of orientational long-range order in their atomic arrangement, and have been studied extensively as potential structural materials due to a unique array of properties such as ultrahigh yield strength, large elastic limit and excellent corrosion resistance compared to crystalline metals [1,2]. One way to produce amorphous alloys is by rapid quenching of the melt to freeze the liquid structure [3], and bulk metallic glasses are usually produced by this process [4]. The other way of obtaining amorphous materials is to induce crystalline-toamorphous transformation (CTAT) by introduction of lattice defects using plastic deformation [5]. Severe plastic deformation (SPD) as effective method to fabricate bulk nanostructure materials has been drawing considerable attention within the last several decades [6]. Recently, amorphous structures have been obtained via CTAT in TiNi alloys by different SPD techniques, such as shot peening [7], cold-rolling [8], high pressure torsion (HPT) [9] and cold drawing [10]. The HPT technique, deforming a disc-shape sample by torsion under high quasi-hydrostatic pressure, can
* Corresponding author. Structural Materials Unit, National Institute for Materials Science, Sengen 1-2-1, Tsukuba, Ibaraki 305-0047, Japan. Tel.: þ81 80 4065 1571; fax: þ81 298592101. E-mail address: [email protected] (F.Q. Meng). 0966-9795/$ e see front matter Ó 2013 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.intermet.2013.01.021
apply extremely large plastic straining by simply increasing the rotation number [6,11], thereby it is suitable to deform brittle alloys and compounds [11,12]. It was found that some alloys with high glass forming ability (GFA) such as ZreTi, ZreCu and ZreAleNieCu have been partially amorphized from solid state by HPT and accumulative roll bonding (ARB) [13e15], but previous research on the amorphization of ZreCu based alloys mostly were from pure elements. The amorphization from intermetallic compounds in ZreCu based alloys has rarely been reported so far. Zr50Cu50xAlx system is well known to be favorable for glass forming by rapid quenching [16]. In the case of rapid quenching, the GFA of these alloys is strongly dependent on the Al addition, and the maximum DTx ( ¼ Tx e Tg, Tg: glass transition temperature, Tx: onset temperature of crystallization) of 80 K can be obtained in the composition of Zr50Cu40Al10 [17]. However, the effect of Al on the amorphization forming ability in the solid state is still unclear. Compared with rapid quenched amorphous structure, the amorphous structure obtained via CTAT is characterized by microstructural heterogeneity, such as high volume fraction of free volume and nanograins, which is regarded as a potential method to improve the ductility of amorphous materials [18,19]. In the present study, the HPT technique was utilized to induce amorphization in crystallized Zr50Cu50xAlx alloys with multi-phase structures. Based on microstructural investigations and thermal analysis, the composition dependence of the amorphous forming ability via CTAT and mechanism were discussed.
F.Q. Meng et al. / Intermetallics 37 (2013) 52e58
2. Experimental Master alloy ingots of Zr50Cu46Al4, Zr50Cu4Al6 and Zr50Cu40Al10were prepared by arc-melting pure Zr, Cu and Al elements in an argon atmosphere. Hereafter, the sample with 4Al, 6Al and 10Al will be referred to as Zr50Cu46Al4, Zr50Cu4Al6 and Zr50Cu40Al10, respectively. To maintain a low-oxygen concentration in the alloys, a high purity Zr with oxygen content less than 0.05 at% was used. The master ingots of alloys were then completely remelted as least three times and then casted into a cylindrical rods with diameter of 10 mm by tilt-casting method in an arc furnace. The X-ray analysis of as-cast samples reveals that 6Al and 10Al alloys were composed of the amorphous phase. The crystallization heat treatment was carried out at 1073 K for 72 h followed by water quenching, which is sufficient to crystallize the amorphous in 6Al and 10Al. The discs for HPT experiment were sliced from crystallized rods with thickness of 0.85 mm. The discs were deformed at room temperature (291 K) using high-pressure torsion apparatus with pressure of 5 GPa and rotation speed of 1 rpm (revolution per minute) for 1, 10, 20 and 50 revolutions. The temperature measured by a thermo couple embedded in the upper anvil after compression and rotation with 1, 10, 20 and 50 revolutions was 291 K, 293 K, 308 K, 314 K and 323 K, respectively. Both planar and cross sections of the deformed samples were prepared; the planar samples were mechanically ground down to roughly the median plane. X-ray diffractometry (XRD) was performed on the planar samples using RINT TTRIII with a Cu-Ka radiation (40 kV, 300 mA). The cross sections of the plastically deformed disc samples were prepared for optical microscopic (OM) and scanning electron microscopic (SEM) observations; the samples were mechanically polished to mirror-like surface using SiO2 colloidal suspension with 0.06 mm in final stage of preparation. OM and SEM observations were performed on a Nikon Eclipse LV150 microscope in polarized light mode, and JEOL 7001F operated at 20 kV
53
in the backscatter electron (BSE) mode, respectively. Disc samples (f3 mm) for TEM observations were cut at positions 3.5 mm from the center of the deformed samples, and then mechanically thinned to 150 mm thick and perforated by electropolishing using a Tenupol5 with an electrolyte consisting of 20 vol% nitric acid and methanol at a temperature of 253 K. TEM observations were performed on JEOL JEM-2000FX and 2100F microscopes with an accelerating voltage of 200 kV. Differential scanning calorimetry (DSC) was carried out in a PerkineElmer Diamond DSC with a heating rate of 6.67 K/s. 1 mm 1 mm 0.8 mm samples with the mass of w30 mg for DSC measurements were cut from the edge of the HPT disc at the same radius as the preparation for TEM sample. 3. Results The XRD profiles of the crystallized 4Al, 6Al and 10Al samples before HPT deformation are shown in Fig. 1 (a). It can be seen clearly that all the samples were in the fully crystallized states. In the 4Al sample, two kinds of martensite phases are detected, one with a base structure (B190 ) with P21/m symmetry and the other with its superstructure with Cm symmetry structure [20]. Most of the peaks for B190 and Cm martensite phases overlap and there are only a few minor peaks that can be solely attributed to B190 at around 42 [21]. The corresponding OM microstructure of 4Al is characterized by coarse martensite plates as shown in Fig. 1 (b). The greater Al addition in 6Al and 10Al alloys leads to a complex multiphase configuration, and Zr2Cu, s5, ZrCu with B2 structure are detected under XRD analysis besides B190 and Cm martensite phases, which is in agreement with the phases diagram. From previous reports [21], it is shown that in increasing Al content from 4Al, 6Al and 10Al leads to the decrease in Ms (forward martensitic transformation starting temperature) at 362 K, 311 K and 306 K and As (reverse martensitic transformation starting temperature) at
Fig. 1. (a) XRD profiles of 4Al, 6Al and 10Al before HPT deformation. (b) Optical micrograph of 4Al. (c) and (d) BSE-SEM images of 6Al and 10Al. Zr2Cu and s5 are marked by the black and white arrows in (c), respectively.
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474 K, 464 K and 457 K, respectively. The decrease in Ms also hinders the transformation from B2 to martensite phases, which explains the presence of ZrCu phase with B2 structure at room temperature. The BSE-SEM images of the 6Al and 10Al samples are shown in Fig. 1 (c) and (d). The contrast difference in the BSE-SEM images (Fig. 1(c) and (d)) reveals the compositional change in different phases. It can be seen clearly that the Al concentrates in s5 phase indicated by bright arrow, and Zr concentrates in Zr2Cu phase indicated by black arrow, as shown in Fig. 1 (c). The volume fraction of Zr2Cu and s5 phases in 10Al is much higher than them in 6Al, which is consistent with the phase diagram [22]. The XRD patterns measured on the planar sections of the HPT-deformed samples with different rotation number are shown in Fig. 2, in which the N denotes to the HPT rotation number and N ¼ 0 represents the sample deformed by only compression without rotation. It can be seen clearly that the diffraction peaks are weakened and broadened gradually with increasing rotation number in 4Al, 6Al and 10Al samples, indicating the refinement of crystalline size and increase of lattice strain after HPT. Neither significant shifts of XRD peaks nor additional diffraction peak is detected during HPT deformation, which indicates there are no new phases formed. After N ¼ 10, a broad halo at round 2q ¼ 38 begins to appear on the curves with some remaining crystalline diffraction peaks, which implies the formation of amorphous phase. No strong diffraction peaks but only a broad hump can be found in the 4Al and 10Al samples after HPT deformation with N ¼ 50, but it appears only after N ¼ 20 in 6Al, which suggests that the HPT deformation leads to the transformation from crystalline to amorphous. Simultaneously, diffraction peaks from B2 peaks in 6Al and 10Al cannot be detected after torsional straining, which is probably caused by
the stress-induced martensitic transformation from B2 to Cm and B190 phases [23]. In order to clarify the microstructural evolution, BSE-SEM and TEM observations were conducted on the cross sections after HPT deformation. Fig. 3 shows the BSE-SEM microstructure evolution dependence on the rotation number in the 10Al sample. The observation was performed at the position about 4e4.5 mm off the center along the radius in the each sample. The contrast difference in the BSE-SEM images corresponds to the chemical composition change in the different phases and can be used to label the microstructure evolution during HPT deformation. No visible change is found between the sample after compression without rotation (Fig.3 (a)) and as-crystallized one (Fig. 1 (d)), but severe plastic deformation has already achieved in the sample with N ¼ 10 (Fig. 3(b)). It can be seen that the deformation localizes in the region with river flow shape, and the surrounding part has also been distorted strongly. After HPT deformation with N ¼ 20, the sample has been deformed severely. The HPT deformation induces a decrease in the size and volume fraction of bright phase (Zr2Cu), and the grain size of black phase (s5) after 20 turns compared with the as-crystallized one. The consecutive rotation up to N ¼ 50 induces disappearance of the phase separation after HPT deformation as shown in Fig. 3 (d), which suggests that a large extent of atomic mixing of Zr, Cu and Al elements has taken place, in a similar way to mechanical alloying. The bright-field (BF) TEM images of 4Al before and after HPT deformation are shown in Fig. 4. Before HPT deformation, only martensite phases with twins are present in the sample, which is consistent with the OM results. After HPT deformation with N ¼ 10, the original structure with sharp twins disappears and the spots of SAED pattern also are elongated as shown in the inset B of Fig. 4 (b). One shear band crossing the image is found in the sample with
Fig. 2. XRD profiles of (a) 4Al, (b) 6Al and (c) 10Al samples after various HPT rotations.
F.Q. Meng et al. / Intermetallics 37 (2013) 52e58
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Fig. 3. BSE-SEM images of the cross sections in 10Al sample after various HPT rotation (a) N ¼ 0, (b) N ¼ 10, (c) N ¼ 20 and N ¼ 50.
N ¼ 10. Amorphization has been accomplished inside of shear band as indicated by the SAED pattern in the inset A of Fig. 4 (b). With further HPT rotation, the volume fraction of twin structure decreases. Fragmented martensite phase surrounded by amorphous or nanocrystalline structure can be observed in the sample with N ¼ 20. As rotation number up to 50, twin structure cannot be observed under TEM, while a small volume fraction of nanograins distributes in the amorphous matrix, which can be observed in the image and also indicated by the scattered spots in the SAED pattern Fig. 4 (d).
The microstructural evolution in 6Al is also shown in Fig. 5. It can be seen that martensite phases with twin structure are present in the as-crystallized sample, which is as same as it the 4Al. After HPT deformation with N ¼ 10 as shown in Fig. 4 (b), the spots of SAED pattern are also elongated, which implies the severe plastic deformation accumulated in the martensite phase. An amount of twins and dislocations can also be observed in the s5 and Zr2Cu phases indicated in the area A and C (Fig. 5 (b)), respectively. Microstructural observation in the sample deformed with N ¼ 20 indicates the formation of amorphous phase after 20 revolutions,
Fig. 4. Bright-field TEM images of 4Al sample: (a) before HPT deformation with the insert SAED pattern, (b) N ¼ 10 with the inset patterns (A) and (B) obtained from the A and B, respectively. (c) N ¼ 20 with the patterns. (d) N ¼ 50 with the inset patterns.
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Fig. 5. Bright-field TEM images of the 6Al: (a) before HPT deformation with the insert SAED pattern, (b) N ¼ 10 with the inset patterns (A), (B) and (c) obtained from the A, B and C, respectively. (c) N ¼ 20 and (d) N ¼ 50 with the inset patterns.
which can also be confirmed by the SAED pattern in Fig. 5 (c). While there are still some grains with different sizes in this sample, the coarse grains in the bottom right of Fig. 5 (c) indicates the structural inhomogeneity. Compared to the 4Al at N ¼ 20, the amorphization in 6Al has been accomplished at N ¼ 20. The consecutive HPT rotation up to N ¼ 50 induces the more perfect diffraction halo obtained in the 6Al, while an amount of nano-grains with size less than w100 nm coexists in the amorphous structure. In the 10Al sample, small volume fraction of martensite phases is surrounded by the Zr2Cu and s5 phases as shown in Fig. 6 (a). After HPT deformation with N ¼ 10 (Fig. 6 (b)), the martensite
phases cannot be observed, indicating the amorphization has been achieved in the twin structure, but the coarse grains corresponding to Zr2Cu and s5 phases still exist. The SAED pattern from area A reveals the amorphous structure. The elongated spots in SAED pattern obtained from area B indicate that deformation has been accumulated in phase s5 phase (Fig. 6 (b)), associating with the twins and dislocation inside of the grains. The volume fraction of amorphous phase continuously increases with the consecutive HPT deformation as N ¼ 20 and the nanograins with size 20e100 nm are observed instead of coarse grains. The image and SAED pattern obtained in the N ¼ 50 sample also indicate the uniform
Fig. 6. Bright-field TEM images of the 4Al sample: (a) before HPT deformation, (b) N ¼ 10 with the insert SAED patterns (A) and (B) obtained from the A and B, respectively. (c) N ¼ 20 and (d) N ¼ 50 with the inset patterns.
F.Q. Meng et al. / Intermetallics 37 (2013) 52e58
combination of amorphous and nanocrystalline structure is obtained, which is as same as in 4Al and 6Al. In addition, the size and distribution of nanograins in Fig. 6 (d) become more homogeneous in comparison with Fig. 6 (c). DSC measurements were conducted to clarify the thermal behaviors of the amorphous structure obtained in the 4Al, 6Al and 10Al samples with N ¼ 50 and the results are shown in Fig. 7. The existence of an endothermic glass transition prior to a larger exothermic crystallization peak can be clearly seen in 4Al, 6Al and 10Al. The presence of a glass transition temperature (Tg) is strong evidence for the intrinsic character of short-range and medium-range atomic structures in the amorphous phase obtained by HPT, which is similar to the as-cast metallic glass. Almost identical glass transition temperatures in the amorphous structure prepared by rapid quenching (RQ) [24] and HPT technique as shown in Table 1 indicate the chemical homogeneity in the amorphous structure of 4Al, 6Al and 10Al. Below the glass transition temperature, there is a broad exothermic reaction starting at w470 K, which is attributed to the structural relaxation of the amorphous sample. This relaxation exothermic peak is usually found in bulk metallic glasses deformed by surface plastic deformation [25], cold-rolling [26] or HPT [12]. In the present study, it is believed that annihilation of free volume and residual stress in the deformed samples lead to the structural relaxation at lower temperature before glass transition, but the growth of nanograins is depressed during relaxation, which has been confirmed by TEM observation (not shown). 4. Discussion The present results show unambiguously that high pressure torsion can lead to almost full amorphization in the Zr50Cu50xAlx (x ¼ 4, 6, 10) samples. Consistent with previous work using pure elements reported in CueZreAl deformed by HPT [27], and CueZre Ti multilayer deformed by cold-rolling [14], we showed that almost completely amorphous structure can be obtained from the intermetallic compounds by HPT and larger shear strain results in higher amount of amorphous phases. The TEM observations indicate that ease of amorphization in these three alloys is dependent not only on the shear strain, but also on the chemical composition: the
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Table 1 Comparison of glass transition temperatures (Tg) in the amorphous structure obtained by HPT and rapid quenching (RQ) [23].
HPT RQ
4Al
6Al
10Al
700 696
704 704
710 709
amorphization is most pronounced in 10Al, intermediate in 6Al and weakest in 4Al, which is almost the same tendency as the glassforming ability in alloys prepared by rapid cooling [17]. The driving force for the transformation from crystalline to amorphous phases is considered to be the free energy difference between these two structures [28]. In present study, HPT deformation generates large defect densities, especially of dislocations, vacancies, grain and sub-grain boundaries (due to refinement of crystalline size). The higher torsional strain creates more crystalline defects which increase the free energy and eventually induces the destabilization of the crystalline phases when the density of defects reaches the critical value [29]. Additionally, the twin and multiphase boundaries provide an amount of preferential sites for the accumulation of crystalline defects. The preferential amorphization in the martensite phase in 10Al with N ¼ 10 also illustrated the twin boundaries play an important role in CTAT process. High Al content increases the amount of Zr2Cu and s5, which seems to result in the pronounced amorphization via mechanical alloying in 6Al and 10Al. On the other hand, the facilitation of amorphization is related to lower Ms in 6Al (311 K) and 10Al (306 K). The temperature rise during HPT processing already exceeded the Ms in 6Al and 10Al after 20 and 10 revolutions but well below the As, respectively. The lattice structure becomes unstable when the temperature is higher than Ms, facilitating the transformation from crystalline (with the high defect density) to amorphous when the density of crystal defects reaches the critical level [29]. However, the absence of B2 phase by XRD analysis and TEM observation in the deformed samples reveals the amorphization occurs in martensite phase, but not B2 phase. Amorphization progress in the multi-phase structures also requires intermixing, such as in Al sample. The large negative heat of mixing, OHmix, between the alloying elements also plays an important role on the amorphization progress. In ZreCueAl systems, there is a large negative OHmix between Zr and (Cu, Al) and the other factor is the mutual atomic diffusion [13]. The homogeneous chemical distribution after HPT deformation with N ¼ 50 in the BSE-SEM images (Fig. 3 (d)) also indicates the amorphization progress is accompanied by atomic inter-diffusion. It is reported [13] that the diffusion coefficient D can be enhanced by several orders of magnitude by lattice defects or interfaces introduced into a crystalline sample during heavy deformation. Obviously, plastic straining by HPT deformation accelerates the atomic diffusion by increasing the density lattice defects. Less chemical mixing is required in ZrCu(Al) phases, resulting in the preferential amorphization. 5. Conclusion
Fig. 7. Thermal behavior of 4Al, 6Al and 10Al after HPT deformation with N ¼ 50. Tg denotes to the glass transition temperature.
HPT can induce crystalline to amorphous transformation in Zre CueAl alloys, with an almost fully amorphous structure obtained at after 50 revolutions. The amorphous samples produced by HPT have almost identical glass transformation temperature to samples obtained by rapid quenching. The ability for transformation of crystalline to amorphous structure induced by HPT shows similar chemical composition dependence with the technique of rapid quenching. The martensite phases with twin structure were found
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to contribute significantly to the crystalline to amorphous transformation induced by HPT. Higher amount of Zr2Cu and s5 phases and lower martensitic transformation starting temperature lead to the more pronounced amorphization in Zr50Cu44Al6 and Zr50Cu40Al10. Acknowledgements This work is partly supported by the Grant-in-Aid for Scientific Research on Innovative Area, “Bulk Nanostructured Metals”, through MEXT, Japan (contract No. 22102004) and also by Iketani Science Foundation. The work was conducted under the interuniversity cooperative research program of the Institute for Materials Research, Tohoku University. References [1] [2] [3] [4] [5] [6] [7]
Schuh CA, Hufnagel TC, Ramamurty U. Acta Mater 2007;55:4067e109. Ashby MF, Greer AL. Scripta Mater 2006;54:321e6. Klement W, Willens RH, Duwez P. Nature 1960;187:869e70. Johnson WL. MRS Bull 1999;24:15. Okamoto PR, Lam NQ, Rehn LE. Solid State Phys 1998;52:135. Valiev RZ, Alexandrov IV. Nanostruct Mater 1999;12:35e40. Grant DM, Green SM, Wood JV. Acta Metall Mater 1995;43:1045e51.
[8] Nakayama H, Tsuchiya K, Umemoto M. Scripta Mater 2001;44:1781e5. [9] Waitz T. Acta Mater 2005;53:2273e83. [10] Tsuchiya K, Hada Y, Koyano T, Nakajima K, Ohnuma M, Koike T, et al. Scripta Mater 2009;60:749e52. [11] Zhilyaev AP, Langdon TG. Prog Mater Sci 2008;53:893e979. [12] Meng FQ, Tsuchiya K, Ii S, Yokoyama Y. App Phys Lett 2012;101:121913e4. [13] Sagel A, Sieber H, Fecht HJ, Perepezko JH. Acta Mater 1998;46:4233e41. [14] Dinda GP, Rosner H, Wilde G. J Non-Cryst Solids 2007;353:3777e81. [15] Atzmon M, Unruh KM, Johnson WL. J Appl Phys 1985;58:3865e70. [16] Inoue A, Kawase D, Tsai AP, Zhang T, Masumoto T. Mater Sci Eng A 1994;178: 255e63. [17] Yokoyama Y, Fukaura K, Inoue A. Mater Sci Eng A 2004;375-377:427e31. [18] Pauly S, Gorantla S, Wang G, Kuhn U. J Eckert, Nat Mater 2010;9:473e7. [19] Pauly S, Das J, Bednarcik J, Mattern N, Kim KB, Kim DH, et al. Scripta Mater 2009;60:431e4. [20] Koval YN, Firstov GS, Kotko AV. Scripta Metall Mater 1992;27:1611e6. [21] Meng FQ, Tsuchiya K, Yin FX, Ii S, Yokoyama Y. J Alloy Compd 2012;522: 136e40. [22] Markiv VY, Burnashova VV. Poroshk Metall 1970;12:53e8. [23] Matsumoto O, Miyazaki S, Otsuka K, Tamura H. Acta Metall 1987;35:2137e44. [24] Yamada M, Tanimoto Y, Yamasaki T, Kikuchi T, Yokoyama Y, Inoue A. J Soc Mater Sci Jpn 2010;59:124e9. [25] Tian JW, Shaw LL, Wang YD, Yokoyama Y, Liaw PK. Intermetallics 2009;17: 951e7. [26] Liu JW, Cao QP, Chen LY, Wang XD, Jiang JZ. Acta Mater 2010;58:4827e40. [27] Sun YF, Nakamura T, Todaka Y, Umemoto M, Tsuji N. Intermetallics 2009;17: 256e61. [28] Kuhlmann-wilsdorf D, Bednar MS. Scripta Metall Mater 1993;28:371e6. [29] Koike J, Parkin DM, Nastasi M. J Mater Res 1990;5:1414e8.
Contents lists available at SciVerse ScienceDirect
Intermetallics journal homepage: www.elsevier.com/locate/intermet
Crystalline to amorphous transformation in ZreCueAl alloys induced by high pressure torsion F.Q. Meng a, b, *, K. Tsuchiya a, b, Y. Yokoyama c a
Graduate School of Pure and Applied Sciences, University of Tsukuba, Tsukuba, Ibaraki 305-8577, Japan Structural Materials Unit, National Institute for Materials Science, Tsukuba, Ibaraki 305-0047, Japan c Institute for Materials Research, Tohoku University, Sendai, Miyagi 980-8577, Japan b
a r t i c l e i n f o
a b s t r a c t
Article history: Received 27 November 2012 Received in revised form 15 January 2013 Accepted 22 January 2013 Available online 26 February 2013
ZreCueAl alloys with various Al contents were processed by high pressure torsion (HPT). It was found that HPT can induce a remarkable crystalline to amorphous transformation in these alloys, with an almost fully amorphous structure obtained after 50 revolutions. Microstructural investigations indicated that the transformation from crystalline to amorphous structure is strongly dependent on both the accumulative torsional strain and chemical composition. Higher accumulative shear strain and Al content led to more pronounced amorphization. The amorphization occurred preferentially in the martensite phases with twin structure in Zr50Cu40Al10. Differential scanning calorimetric (DSC) analysis revealed a similar glass transition temperature in the amorphous samples produced by HPT to those prepared by rapid quenching. Ó 2013 Elsevier Ltd. All rights reserved.
Keywords: A. Multiphase intermetallics B. Glasses, metallic D. Microstructure D. Martensitic structure
1. Introduction Amorphous alloys are characterized by the absence of orientational long-range order in their atomic arrangement, and have been studied extensively as potential structural materials due to a unique array of properties such as ultrahigh yield strength, large elastic limit and excellent corrosion resistance compared to crystalline metals [1,2]. One way to produce amorphous alloys is by rapid quenching of the melt to freeze the liquid structure [3], and bulk metallic glasses are usually produced by this process [4]. The other way of obtaining amorphous materials is to induce crystalline-toamorphous transformation (CTAT) by introduction of lattice defects using plastic deformation [5]. Severe plastic deformation (SPD) as effective method to fabricate bulk nanostructure materials has been drawing considerable attention within the last several decades [6]. Recently, amorphous structures have been obtained via CTAT in TiNi alloys by different SPD techniques, such as shot peening [7], cold-rolling [8], high pressure torsion (HPT) [9] and cold drawing [10]. The HPT technique, deforming a disc-shape sample by torsion under high quasi-hydrostatic pressure, can
* Corresponding author. Structural Materials Unit, National Institute for Materials Science, Sengen 1-2-1, Tsukuba, Ibaraki 305-0047, Japan. Tel.: þ81 80 4065 1571; fax: þ81 298592101. E-mail address: [email protected] (F.Q. Meng). 0966-9795/$ e see front matter Ó 2013 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.intermet.2013.01.021
apply extremely large plastic straining by simply increasing the rotation number [6,11], thereby it is suitable to deform brittle alloys and compounds [11,12]. It was found that some alloys with high glass forming ability (GFA) such as ZreTi, ZreCu and ZreAleNieCu have been partially amorphized from solid state by HPT and accumulative roll bonding (ARB) [13e15], but previous research on the amorphization of ZreCu based alloys mostly were from pure elements. The amorphization from intermetallic compounds in ZreCu based alloys has rarely been reported so far. Zr50Cu50xAlx system is well known to be favorable for glass forming by rapid quenching [16]. In the case of rapid quenching, the GFA of these alloys is strongly dependent on the Al addition, and the maximum DTx ( ¼ Tx e Tg, Tg: glass transition temperature, Tx: onset temperature of crystallization) of 80 K can be obtained in the composition of Zr50Cu40Al10 [17]. However, the effect of Al on the amorphization forming ability in the solid state is still unclear. Compared with rapid quenched amorphous structure, the amorphous structure obtained via CTAT is characterized by microstructural heterogeneity, such as high volume fraction of free volume and nanograins, which is regarded as a potential method to improve the ductility of amorphous materials [18,19]. In the present study, the HPT technique was utilized to induce amorphization in crystallized Zr50Cu50xAlx alloys with multi-phase structures. Based on microstructural investigations and thermal analysis, the composition dependence of the amorphous forming ability via CTAT and mechanism were discussed.
F.Q. Meng et al. / Intermetallics 37 (2013) 52e58
2. Experimental Master alloy ingots of Zr50Cu46Al4, Zr50Cu4Al6 and Zr50Cu40Al10were prepared by arc-melting pure Zr, Cu and Al elements in an argon atmosphere. Hereafter, the sample with 4Al, 6Al and 10Al will be referred to as Zr50Cu46Al4, Zr50Cu4Al6 and Zr50Cu40Al10, respectively. To maintain a low-oxygen concentration in the alloys, a high purity Zr with oxygen content less than 0.05 at% was used. The master ingots of alloys were then completely remelted as least three times and then casted into a cylindrical rods with diameter of 10 mm by tilt-casting method in an arc furnace. The X-ray analysis of as-cast samples reveals that 6Al and 10Al alloys were composed of the amorphous phase. The crystallization heat treatment was carried out at 1073 K for 72 h followed by water quenching, which is sufficient to crystallize the amorphous in 6Al and 10Al. The discs for HPT experiment were sliced from crystallized rods with thickness of 0.85 mm. The discs were deformed at room temperature (291 K) using high-pressure torsion apparatus with pressure of 5 GPa and rotation speed of 1 rpm (revolution per minute) for 1, 10, 20 and 50 revolutions. The temperature measured by a thermo couple embedded in the upper anvil after compression and rotation with 1, 10, 20 and 50 revolutions was 291 K, 293 K, 308 K, 314 K and 323 K, respectively. Both planar and cross sections of the deformed samples were prepared; the planar samples were mechanically ground down to roughly the median plane. X-ray diffractometry (XRD) was performed on the planar samples using RINT TTRIII with a Cu-Ka radiation (40 kV, 300 mA). The cross sections of the plastically deformed disc samples were prepared for optical microscopic (OM) and scanning electron microscopic (SEM) observations; the samples were mechanically polished to mirror-like surface using SiO2 colloidal suspension with 0.06 mm in final stage of preparation. OM and SEM observations were performed on a Nikon Eclipse LV150 microscope in polarized light mode, and JEOL 7001F operated at 20 kV
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in the backscatter electron (BSE) mode, respectively. Disc samples (f3 mm) for TEM observations were cut at positions 3.5 mm from the center of the deformed samples, and then mechanically thinned to 150 mm thick and perforated by electropolishing using a Tenupol5 with an electrolyte consisting of 20 vol% nitric acid and methanol at a temperature of 253 K. TEM observations were performed on JEOL JEM-2000FX and 2100F microscopes with an accelerating voltage of 200 kV. Differential scanning calorimetry (DSC) was carried out in a PerkineElmer Diamond DSC with a heating rate of 6.67 K/s. 1 mm 1 mm 0.8 mm samples with the mass of w30 mg for DSC measurements were cut from the edge of the HPT disc at the same radius as the preparation for TEM sample. 3. Results The XRD profiles of the crystallized 4Al, 6Al and 10Al samples before HPT deformation are shown in Fig. 1 (a). It can be seen clearly that all the samples were in the fully crystallized states. In the 4Al sample, two kinds of martensite phases are detected, one with a base structure (B190 ) with P21/m symmetry and the other with its superstructure with Cm symmetry structure [20]. Most of the peaks for B190 and Cm martensite phases overlap and there are only a few minor peaks that can be solely attributed to B190 at around 42 [21]. The corresponding OM microstructure of 4Al is characterized by coarse martensite plates as shown in Fig. 1 (b). The greater Al addition in 6Al and 10Al alloys leads to a complex multiphase configuration, and Zr2Cu, s5, ZrCu with B2 structure are detected under XRD analysis besides B190 and Cm martensite phases, which is in agreement with the phases diagram. From previous reports [21], it is shown that in increasing Al content from 4Al, 6Al and 10Al leads to the decrease in Ms (forward martensitic transformation starting temperature) at 362 K, 311 K and 306 K and As (reverse martensitic transformation starting temperature) at
Fig. 1. (a) XRD profiles of 4Al, 6Al and 10Al before HPT deformation. (b) Optical micrograph of 4Al. (c) and (d) BSE-SEM images of 6Al and 10Al. Zr2Cu and s5 are marked by the black and white arrows in (c), respectively.
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474 K, 464 K and 457 K, respectively. The decrease in Ms also hinders the transformation from B2 to martensite phases, which explains the presence of ZrCu phase with B2 structure at room temperature. The BSE-SEM images of the 6Al and 10Al samples are shown in Fig. 1 (c) and (d). The contrast difference in the BSE-SEM images (Fig. 1(c) and (d)) reveals the compositional change in different phases. It can be seen clearly that the Al concentrates in s5 phase indicated by bright arrow, and Zr concentrates in Zr2Cu phase indicated by black arrow, as shown in Fig. 1 (c). The volume fraction of Zr2Cu and s5 phases in 10Al is much higher than them in 6Al, which is consistent with the phase diagram [22]. The XRD patterns measured on the planar sections of the HPT-deformed samples with different rotation number are shown in Fig. 2, in which the N denotes to the HPT rotation number and N ¼ 0 represents the sample deformed by only compression without rotation. It can be seen clearly that the diffraction peaks are weakened and broadened gradually with increasing rotation number in 4Al, 6Al and 10Al samples, indicating the refinement of crystalline size and increase of lattice strain after HPT. Neither significant shifts of XRD peaks nor additional diffraction peak is detected during HPT deformation, which indicates there are no new phases formed. After N ¼ 10, a broad halo at round 2q ¼ 38 begins to appear on the curves with some remaining crystalline diffraction peaks, which implies the formation of amorphous phase. No strong diffraction peaks but only a broad hump can be found in the 4Al and 10Al samples after HPT deformation with N ¼ 50, but it appears only after N ¼ 20 in 6Al, which suggests that the HPT deformation leads to the transformation from crystalline to amorphous. Simultaneously, diffraction peaks from B2 peaks in 6Al and 10Al cannot be detected after torsional straining, which is probably caused by
the stress-induced martensitic transformation from B2 to Cm and B190 phases [23]. In order to clarify the microstructural evolution, BSE-SEM and TEM observations were conducted on the cross sections after HPT deformation. Fig. 3 shows the BSE-SEM microstructure evolution dependence on the rotation number in the 10Al sample. The observation was performed at the position about 4e4.5 mm off the center along the radius in the each sample. The contrast difference in the BSE-SEM images corresponds to the chemical composition change in the different phases and can be used to label the microstructure evolution during HPT deformation. No visible change is found between the sample after compression without rotation (Fig.3 (a)) and as-crystallized one (Fig. 1 (d)), but severe plastic deformation has already achieved in the sample with N ¼ 10 (Fig. 3(b)). It can be seen that the deformation localizes in the region with river flow shape, and the surrounding part has also been distorted strongly. After HPT deformation with N ¼ 20, the sample has been deformed severely. The HPT deformation induces a decrease in the size and volume fraction of bright phase (Zr2Cu), and the grain size of black phase (s5) after 20 turns compared with the as-crystallized one. The consecutive rotation up to N ¼ 50 induces disappearance of the phase separation after HPT deformation as shown in Fig. 3 (d), which suggests that a large extent of atomic mixing of Zr, Cu and Al elements has taken place, in a similar way to mechanical alloying. The bright-field (BF) TEM images of 4Al before and after HPT deformation are shown in Fig. 4. Before HPT deformation, only martensite phases with twins are present in the sample, which is consistent with the OM results. After HPT deformation with N ¼ 10, the original structure with sharp twins disappears and the spots of SAED pattern also are elongated as shown in the inset B of Fig. 4 (b). One shear band crossing the image is found in the sample with
Fig. 2. XRD profiles of (a) 4Al, (b) 6Al and (c) 10Al samples after various HPT rotations.
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55
Fig. 3. BSE-SEM images of the cross sections in 10Al sample after various HPT rotation (a) N ¼ 0, (b) N ¼ 10, (c) N ¼ 20 and N ¼ 50.
N ¼ 10. Amorphization has been accomplished inside of shear band as indicated by the SAED pattern in the inset A of Fig. 4 (b). With further HPT rotation, the volume fraction of twin structure decreases. Fragmented martensite phase surrounded by amorphous or nanocrystalline structure can be observed in the sample with N ¼ 20. As rotation number up to 50, twin structure cannot be observed under TEM, while a small volume fraction of nanograins distributes in the amorphous matrix, which can be observed in the image and also indicated by the scattered spots in the SAED pattern Fig. 4 (d).
The microstructural evolution in 6Al is also shown in Fig. 5. It can be seen that martensite phases with twin structure are present in the as-crystallized sample, which is as same as it the 4Al. After HPT deformation with N ¼ 10 as shown in Fig. 4 (b), the spots of SAED pattern are also elongated, which implies the severe plastic deformation accumulated in the martensite phase. An amount of twins and dislocations can also be observed in the s5 and Zr2Cu phases indicated in the area A and C (Fig. 5 (b)), respectively. Microstructural observation in the sample deformed with N ¼ 20 indicates the formation of amorphous phase after 20 revolutions,
Fig. 4. Bright-field TEM images of 4Al sample: (a) before HPT deformation with the insert SAED pattern, (b) N ¼ 10 with the inset patterns (A) and (B) obtained from the A and B, respectively. (c) N ¼ 20 with the patterns. (d) N ¼ 50 with the inset patterns.
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Fig. 5. Bright-field TEM images of the 6Al: (a) before HPT deformation with the insert SAED pattern, (b) N ¼ 10 with the inset patterns (A), (B) and (c) obtained from the A, B and C, respectively. (c) N ¼ 20 and (d) N ¼ 50 with the inset patterns.
which can also be confirmed by the SAED pattern in Fig. 5 (c). While there are still some grains with different sizes in this sample, the coarse grains in the bottom right of Fig. 5 (c) indicates the structural inhomogeneity. Compared to the 4Al at N ¼ 20, the amorphization in 6Al has been accomplished at N ¼ 20. The consecutive HPT rotation up to N ¼ 50 induces the more perfect diffraction halo obtained in the 6Al, while an amount of nano-grains with size less than w100 nm coexists in the amorphous structure. In the 10Al sample, small volume fraction of martensite phases is surrounded by the Zr2Cu and s5 phases as shown in Fig. 6 (a). After HPT deformation with N ¼ 10 (Fig. 6 (b)), the martensite
phases cannot be observed, indicating the amorphization has been achieved in the twin structure, but the coarse grains corresponding to Zr2Cu and s5 phases still exist. The SAED pattern from area A reveals the amorphous structure. The elongated spots in SAED pattern obtained from area B indicate that deformation has been accumulated in phase s5 phase (Fig. 6 (b)), associating with the twins and dislocation inside of the grains. The volume fraction of amorphous phase continuously increases with the consecutive HPT deformation as N ¼ 20 and the nanograins with size 20e100 nm are observed instead of coarse grains. The image and SAED pattern obtained in the N ¼ 50 sample also indicate the uniform
Fig. 6. Bright-field TEM images of the 4Al sample: (a) before HPT deformation, (b) N ¼ 10 with the insert SAED patterns (A) and (B) obtained from the A and B, respectively. (c) N ¼ 20 and (d) N ¼ 50 with the inset patterns.
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combination of amorphous and nanocrystalline structure is obtained, which is as same as in 4Al and 6Al. In addition, the size and distribution of nanograins in Fig. 6 (d) become more homogeneous in comparison with Fig. 6 (c). DSC measurements were conducted to clarify the thermal behaviors of the amorphous structure obtained in the 4Al, 6Al and 10Al samples with N ¼ 50 and the results are shown in Fig. 7. The existence of an endothermic glass transition prior to a larger exothermic crystallization peak can be clearly seen in 4Al, 6Al and 10Al. The presence of a glass transition temperature (Tg) is strong evidence for the intrinsic character of short-range and medium-range atomic structures in the amorphous phase obtained by HPT, which is similar to the as-cast metallic glass. Almost identical glass transition temperatures in the amorphous structure prepared by rapid quenching (RQ) [24] and HPT technique as shown in Table 1 indicate the chemical homogeneity in the amorphous structure of 4Al, 6Al and 10Al. Below the glass transition temperature, there is a broad exothermic reaction starting at w470 K, which is attributed to the structural relaxation of the amorphous sample. This relaxation exothermic peak is usually found in bulk metallic glasses deformed by surface plastic deformation [25], cold-rolling [26] or HPT [12]. In the present study, it is believed that annihilation of free volume and residual stress in the deformed samples lead to the structural relaxation at lower temperature before glass transition, but the growth of nanograins is depressed during relaxation, which has been confirmed by TEM observation (not shown). 4. Discussion The present results show unambiguously that high pressure torsion can lead to almost full amorphization in the Zr50Cu50xAlx (x ¼ 4, 6, 10) samples. Consistent with previous work using pure elements reported in CueZreAl deformed by HPT [27], and CueZre Ti multilayer deformed by cold-rolling [14], we showed that almost completely amorphous structure can be obtained from the intermetallic compounds by HPT and larger shear strain results in higher amount of amorphous phases. The TEM observations indicate that ease of amorphization in these three alloys is dependent not only on the shear strain, but also on the chemical composition: the
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Table 1 Comparison of glass transition temperatures (Tg) in the amorphous structure obtained by HPT and rapid quenching (RQ) [23].
HPT RQ
4Al
6Al
10Al
700 696
704 704
710 709
amorphization is most pronounced in 10Al, intermediate in 6Al and weakest in 4Al, which is almost the same tendency as the glassforming ability in alloys prepared by rapid cooling [17]. The driving force for the transformation from crystalline to amorphous phases is considered to be the free energy difference between these two structures [28]. In present study, HPT deformation generates large defect densities, especially of dislocations, vacancies, grain and sub-grain boundaries (due to refinement of crystalline size). The higher torsional strain creates more crystalline defects which increase the free energy and eventually induces the destabilization of the crystalline phases when the density of defects reaches the critical value [29]. Additionally, the twin and multiphase boundaries provide an amount of preferential sites for the accumulation of crystalline defects. The preferential amorphization in the martensite phase in 10Al with N ¼ 10 also illustrated the twin boundaries play an important role in CTAT process. High Al content increases the amount of Zr2Cu and s5, which seems to result in the pronounced amorphization via mechanical alloying in 6Al and 10Al. On the other hand, the facilitation of amorphization is related to lower Ms in 6Al (311 K) and 10Al (306 K). The temperature rise during HPT processing already exceeded the Ms in 6Al and 10Al after 20 and 10 revolutions but well below the As, respectively. The lattice structure becomes unstable when the temperature is higher than Ms, facilitating the transformation from crystalline (with the high defect density) to amorphous when the density of crystal defects reaches the critical level [29]. However, the absence of B2 phase by XRD analysis and TEM observation in the deformed samples reveals the amorphization occurs in martensite phase, but not B2 phase. Amorphization progress in the multi-phase structures also requires intermixing, such as in Al sample. The large negative heat of mixing, OHmix, between the alloying elements also plays an important role on the amorphization progress. In ZreCueAl systems, there is a large negative OHmix between Zr and (Cu, Al) and the other factor is the mutual atomic diffusion [13]. The homogeneous chemical distribution after HPT deformation with N ¼ 50 in the BSE-SEM images (Fig. 3 (d)) also indicates the amorphization progress is accompanied by atomic inter-diffusion. It is reported [13] that the diffusion coefficient D can be enhanced by several orders of magnitude by lattice defects or interfaces introduced into a crystalline sample during heavy deformation. Obviously, plastic straining by HPT deformation accelerates the atomic diffusion by increasing the density lattice defects. Less chemical mixing is required in ZrCu(Al) phases, resulting in the preferential amorphization. 5. Conclusion
Fig. 7. Thermal behavior of 4Al, 6Al and 10Al after HPT deformation with N ¼ 50. Tg denotes to the glass transition temperature.
HPT can induce crystalline to amorphous transformation in Zre CueAl alloys, with an almost fully amorphous structure obtained at after 50 revolutions. The amorphous samples produced by HPT have almost identical glass transformation temperature to samples obtained by rapid quenching. The ability for transformation of crystalline to amorphous structure induced by HPT shows similar chemical composition dependence with the technique of rapid quenching. The martensite phases with twin structure were found
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to contribute significantly to the crystalline to amorphous transformation induced by HPT. Higher amount of Zr2Cu and s5 phases and lower martensitic transformation starting temperature lead to the more pronounced amorphization in Zr50Cu44Al6 and Zr50Cu40Al10. Acknowledgements This work is partly supported by the Grant-in-Aid for Scientific Research on Innovative Area, “Bulk Nanostructured Metals”, through MEXT, Japan (contract No. 22102004) and also by Iketani Science Foundation. The work was conducted under the interuniversity cooperative research program of the Institute for Materials Research, Tohoku University. References [1] [2] [3] [4] [5] [6] [7]
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