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Microhardness and abrasive wear resistance of metallic glasses and nanostructured composite materials
Journal of Non-Crystalline Solids 316 (2003) 96–103 www.elsevier.com/locate/jnoncrysol
Microhardness and abrasive wear resistance of metallic glasses and nanostructured composite materials Thierry Gloriant
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Laboratoire de M etallurgie et Physico-Chimie des Mat eriaux, G.R.C.M. (EA2620) Institut National des Sciences Appliqu ees de Rennes 20, avenue des Buttes de Co€esmes, F-35043 Rennes cedex, France
Abstract The present paper reports on as-cast bulk zirconium-based, lanthanum-based and palladium-based metallic glasses and on melt-spun aluminum-based amorphous ribbons. Their glass forming abilities, nano(quasi)crystallization mechanisms and mechanical properties are discussed through a short review of very recent papers and by the presentation of some experimental results. The tribological aspect such as microhardness and abrasive wear resistance are more particularly investigated for their potential applications technologically useful as coatings. The formation of nanostructured composite materials consisting of nanophase particles embedded in a residual amorphous matrix by appropriate thermal treatments is also studied and the influence of the nanocrystallized volume fraction is evaluated and discussed for both hardness and wear resistance. Ó 2003 Elsevier Science B.V. All rights reserved. PACS: 81.05.K; 61.43.D; 62.25
1. Introduction Since the first success of preparing an amorphous phase in the Au–Si system by rapid solidification in 1960 [1] a great number of metallic amorphous alloys has been produced until now. Before 1990, Fe-, Co- and Ni-based amorphous alloys were formed by using high cooling rates but the resulting sample thickness was limited to less than about 50 lm [2] reducing their potential as engineering materials. However, in 1984, an ex-
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ceptional example, Pd40 Ni40 P20 was discovered as the first bulk metallic glass (BMG) forming alloy with a critical cooling rate as low as about 1 K s1 when processed in B2 O3 [3]. Thus for a long period before 1990, no other amorphous alloys except the Pd–Ni–P and Pt–Ni–P systems [4] were synthesized because of the necessity of high cooling rate above 105 K s1 for glass formation. The discovery of the La- [5] and Zr-based metallic glasses [6,7] in the late 1990s from multicomponents systems has triggered tremendous research activity in this area particularly for their excellent mechanical properties. Thus, a number of bulk amorphous alloys have been produced in Zr–Al–TM [6,8], Zr–Ti– TM–Be [7], Ti–Zr–Al–TM–Be [9], Pd–Cu–Ni–P [10] or in La–Al–TM systems [5] (TM ¼ transition metal) by conventional casting processes. It is to
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be noticed however that the maximum sample thickness reaches about 40 mm maximum for Zrbased alloy [11,12] and 70 mm for Pd-based alloy [13]. In order to produce metallic glasses from the melt, the cooling rate must be sufficiently high to avoid nucleation and growth of crystalline phases in the supercooled liquid region between melting temperature Tm and glass transition Tg . As recently reported by Schneider [14], Turnbull [15], using a simple nucleation theory, showed that the reduced glass transition temperature, Trg ¼ Tg =Tm , is a key parameter that determines whether or not the melt of a given material can form a glass during cooling and concluded that to produce bulk metallic glasses, Trg must be relatively high. On the other hand, in order to develop some of the best bulk glass-forming alloys, Johnson [16] has successfully added suitable components (Cu, Ni) to lower the eutectic temperature in the Zr–Ti–Be system obtaining the family of Zr–Ti–Cu–Ni–Be glasses. These alloys show an exceptional glass forming ability and a critical cooling rate as low as about 1 K s1 such as for Vitreloye, which makes them very interesting materials for engineering applications [7]. With the same objective and in order to optimize the stabilization of metallic supercooled liquid and bulk amorphous alloys, Inoue [17] has proposed the following three empirical rules to achieve bulk glass formation: (1) multicomponent systems consisting of more than three elements; (2) significant difference in atomic size ratios above about 12% among the three main constituent elements; and (3) negative heats of mixing among the three main constituent elements. By choosing the above-described multicomponent alloy systems, it is possible to produce amorphous ribbons by melt-spinning (when a very high cooling rate is required), and bulk amorphous alloys by solidification techniques such as copper mold casting, arc melting suction casting, squeeze casting, high pressure die casting, unidirectional melting or water quenching. The standard procedures are given in detail in [18]. Metallic glasses have also attracted increasing interest as precursors to produce nanophase composite materials because of the appearance of good properties such as mechanical [19,20] or
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magnetic [21,22]. This was particularly investigated for Zr-based BMG and also for Al-based metallic glasses after their first elaboration by rapid solidification (melt spinning) in 1988 [23–25]. These amorphous alloys crystallize into a nanostructured composite material consisting of nanometer-sized crystallites (or quasicrystallites) embedded in an amorphous matrix, often further enhancing mechanical properties above those of the metallic glasses [19,20]. Some of the best alloys concerning the glass forming ability of Al-based amorphous alloys (>85 at.%Al) are those elaborated from Al–TM–RE (TM ¼ transistion metal, RE ¼ rare earth) systems [26,27]. However, the glass forming ability is not high enough to obtain them as BMG by conventional casting and unfortunately a rapid solidification process (>105 K s1 ) is still required up to now [28,29]. The devitrified nanostructure can be formed by the suppression of growth reaction caused by segregation of a solute element with low atomic diffusivity at the nanocrystal/amorphous interface and by the higher thermal stability of the remaining amorphous phase due to the solute enrichment occurring during crystallization. The final nanostructure is strongly dependent on the number of stable homogeneous nucleation sites in the as-quenched glass but it has also been suggested a possible control by a heterogeneous nucleation mechanism by additional elements occurring by formation of clusters or by the presence of quenched-in nuclei [30]. In Zr-based amorphous alloys, it has been reported that an icosahedral quasicrystalline phase is formed in some alloys with a supercooled liquid region and high glass forming ability [31,32]. Recently, K€ oster et al. [33] have reviewed the conditions for icosahedral quasicrystal formation upon annealing in Zr–Al–Ni–Cu bulk metallic glasses and the beneficial influence of alloying elements such as Hf or Y. In similar compositions, addition of Ag or Pd to the Zr–Al–Ni–Cu glassy alloys has been observed to be effective for the increase of the number of nucleation sites and for the suppression of grain growth of the nano icosahedral phase [34]. As previously mentioned, devitrification of Albased amorphous ribbons upon annealing leads to the formation of a nanometer-scale dispersion of
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a-Al crystallites (or nanoquasicrystallites, see [35] for review) in the remaining matrix and the nanocrystallization mechanism occurring by nucleation and growth was intensively studies by several authors [26,27,36]. Recently, it was observed by TEM that the crystallization is preceded by a phase separation of the amorphous matrix (of about 40 nm) and nucleation of a-Al nanocrystals starts preferentially at the interface of the phaseseparated regions [37]. Because of the limited thickness obtained in the preparation of metallic glasses, particularly with Al-based alloys, this paper is focused on the attractive results for tribological applications, which may suggest metallic glasses technologically useful as coating by a combination of high hardness and high abrasive wear resistance.
microscopy (3DAPFIM) in order to obtain chemical distribution maps of the elements (Al, Ni and Sm) at the nanometer scale (see [41] for more details concerning this study). The 3DAPFIM technique (described in Ref. [42]) has been proven to be ideal for studying solute clustering, precipitation, segregation or partitioning in nanocrystallization [42]. For the wear resistance measurements, a microscale abrasion test was used. This involves rotating a steel sphere against a test specimen in the presence of small angular SiC particles. This mode of abrasion is commonly termed Õthree-bodyÕ and is described in more details elsewhere [43–45]. The wear resistance may be calculated from the slopes of the curves representing the wear crater volume measured during the abrasion test following the classical model proposed by Rabinowitz et al. [46] according to
2. Experimental methods Vw ¼ This paper presents experimental results on three BMG (Zr-based, Pd-based and La-based) and two Al-based amorphous ribbons. Bulk amorphous Zr65 Al7:5 Ni10 Cu17:5 (16 mm of diameter) and Pd40 Ni10 Cu30 P20 (40 mm of diameter) compositions were prepared by water quenching using the standard procedure described in detail elsewhere [10,38] and the bulk La60 Al20 Ni10 Co5 Cu5 glass (2 mm of diameter) was prepared by a high-pressure die casting method as presented in [39]. Amorphous Al88 Ni4 Sm8 and Al86 Ni11 Yb3 alloys were produced by melt spinning and the ribbons obtained were typically 3 mm wide and 50 lm thick. The volume fraction of nanocrystalline a-Al transformed during different annealing times (in tubular furnace under argon atmosphere) was evaluated using the enthalpy of formation calculated from the first differential scanning calorimetry peak according to the method proposed in Ref. [40]. The as-quenched amorphous state and the crystallization behavior under annealing were characterized by X-ray diffraction and by transmission and high resolution electron microscopy (TEM and HREM) using a Philips CM30 microscope. The Al88 Ni4 Sm8 partially crystallized alloy was particularly investigated by three-dimensional atom-probe field-ion
SN ; Rw
ð1Þ
where Vw is the wear volume measured with a calibrated optical microscope, S, the sliding distance, N , the normal load and Rw , the wear resistance. Vickers testers were used with a 0.20 or 0.49 N load for the microhardness measurements.
3. Results and discussion 3.1. Hardness and wear resistance of metallic glasses It has been described in Section 1 that the glassy alloy exhibit high strength combined with high toughness leading to interesting practical use as engineering materials. For example, bulk La–Al– TM (TM ¼ transition metal) alloys exhibit high tensile strength [39], Zr-based amorphous alloys present superplastic deformation [47] and high impact fracture energy [48] and Al–Ni–RE (RE ¼ rare earth metal) amorphous ribbons exhibit high tensile strength above 980 MPa and good bending ductility [24]. For these reasons, metallic glasses have been extensively studied but there is not a lot of work concerning the wear resistance.
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There are many possibilities to evaluate the abrasive properties of materials such as pin-on disk, pin-on plate, cylinder, ball or wheel on flat or scratch test. . . The first documented wear test on a metallic glass was by Boswell in 1997 who studied the abrasive wear of Pd78:1 Si61:4 Cu5:5 by using a pin-on-disk set-up [49] but most of the studies in this field were concentrated on iron-based alloys because these were the first to become available in industrial quantities and to present a relatively good wear resistance [50]. By using the technique presented in this paper, Greer et al. were the first to present results concerning Al, Zr, Pd and Labased amorphous alloys [51–53]. Fig. 1 displays the volume of the wear craters (measured as a function of the sliding distance) in metallic glasses representing four common types of glass compositions, i.e. an Al-based amorphous ribbon and bulk amorphous zirconium, lanthanum and palladium based alloys as described in Section 2. The measurements were carried out by interrupted tests as a function of sliding distance and show that Eq. (1) is obeyed by each of the metallic glasses. Fig. 2 compares the wear resistance of the four metallic glasses in correlation with their microhardness values and compared with a wide variety of materials by using the same micro-scale abrasive wear test. It can be seen that the wear resistance in-
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Fig. 2. The correlation of abrasive wear resistance with microhardness for different material classes (measured using the same technique as in the present work). The four glasses are represented by open squares [44,45].
creases linearly with the hardness, as expected within a given material class. This comparison of materials, as observed in Fig. 2, indicates that the metallic glasses fit firmly in the classification of hardened alloys. For a given hardness, hardened alloys show wear resistances significantly greater than those of ceramics (which are reduced by brittle failure mechanisms) and significantly less than those of pure metals (which are enhanced by work hardening). The wear resistance of metallic glasses, much greater than that of the ceramics of the same hardness, suggests that the wear mechanism does not involve brittle fracture, which is of great interest for coating applications. 3.2. Nanocrystallization of metallic glasses
Fig. 1. The volume wear craters in the four metallic glasses measured as a function of sliding distance during interrupted tests.
An example of a devitrified nanostructure is shown in the TEM dark-field image presented in Fig. 3. The micrograph shows the dispersion of aAl nanocrystallites (in bright) of about 20–30 nm in size precipitated after a thermal treatment at 500 K during 120 min of the Al88 Ni4 Sm8 amorphous composition. This composition was also studied by high-resolution electron microscopy (see the HREM micrograph in Fig. 4 showing a nanocrystal/amorphous interface) and by threedimensional atom-probe field-ion microscopy
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partitioning is confirmed to play a significant role in impeding their growth and aiding the formation of a fine-scale microstructure by the rejection of the two solute elements from the pure a-Al crystals to the matrix during crystallization. 3.3. Effect of the nanocrystallization on hardness and wear resistance
Fig. 3. TEM dark-field image of Al88 Ni4 Sm8 amorphous alloy heat-treated at 500 K for 120 min.
Fig. 4. High-resolution transmission electron microscope image (HREM) of Al88 Ni4 Sm8 (annealed: 500 K/120 min).
(3DAPFIM) (results have been reported in a recent paper [41]). As presented in Fig. 5, it is observed by 3DAPFIM that nanocrystallites appear to be completely devoid of solute content (Ni and Sm). The three-dimensional map on the left side of Fig. 5 represents the samarium content in the melt (each point is one Sm atom) and shows a region devoid of Sm atom. This zone corresponds to the 100 at.% aluminium isoconcentration surface presented on the right 3D map in the figure showing one pure a-Al nanocrystal. Consequently, solute
It is well established that mechanical properties are further enhanced by homogeneous dispersion of nanoscale compound particles into the glassy phase. For instance, Inoue et al. [54] have observed that precipitation of nanoscale quasicrystalline particles with a size of about 30 nm in the Zr–Al– Ni–Cu–(Ag or Pd) bulk glassy alloy exhibits enhanced mechanical properties such as higher tensile or compressive strength, larger elongation an higher hardness and Eckert et al. [55] have observed high strength and good ductility in nanostructured Zr-based multiphase alloys. For a La-based BMG, recent experimental results show that introducing nanocrystalline particles by way of heat treating the amorphous alloy can increase its stiffness and strength [56]. On the other hand, Al-based nanophase composites are known to present quite remarkable tensile strength at room temperature as high as about 1.5 times that of the corresponding fully amorphous alloy and about 3 times that of a conventional precipitation-hardened alloy [57]. For tribological interest, it was shown that hardness and wear resistance are also enhanced by nanometer-scale dispersion of crystallites or quasicrystallites in an amorphous matrix. By studying nano-scratch behavior of bulk Zr-based alloys, Wang et al. [58] have recently shown that samples having a structure consisting of an amorphous matrix with dispersed nanocrystalline particles exhibits the best wear performance. In amorphous alloys from Al–Ni–RE systems, hardness and abrasive wear resistance were also enhanced by nanocrystallization [51,59– 61]. Fig. 6 presents experimental results obtained for the melt spun Al86 Ni11 Yb3 amorphous ribbon. This figure reveals a clear correlation between the microhardness and the volume fraction of a-Al nanocrystallites obtained by isothermal annealing at 473 K. The microhardness is shown to increase
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Fig. 5. Three-dimensional Sm atom distribution (by 3DAPFIM) after annealing Al88 Ni4 Sm8 at 500 K for 15 min (left). Corresponding 100 at.% Al isoconcentration surface (right) [28].
Fig. 6. Volume fraction of a-Al nanoparticles crystallized and microhardness of Al86 Ni11 Yb3 alloy for different times of annealing at 473 K [61].
more than twice, reaching about 400 Hv for a crystallized volume fraction of about 25%. The wear resistance is also observed to be enhanced with the a-Al volume fraction crystallized as shown the Fig. 7. In this figure are presented wear resistance and volume fraction curves of the Al88 Ni4 Sm8 alloy isothermally annealed at 500 K.
Fig. 7. Volume fraction of a-Al nanoparticles crystallized and abrasive wear resistance of Al88 Ni4 Sm8 alloy for different times of annealing at 500 K [60].
After about 120 min of heat treatment, the wear resistance has increased from 1.5 to 2.3.10þ11 Pa induced by a nanocrystalline volume fraction of about 40%. The mechanism that can explained the observed increase of mechanical properties such as hardness and wear resistance by formation of nanophase
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particles in a residual amorphous matrix has been extensively discussed. Kim et al. [57] have pointed out that the nanocrystallites are too small to content defects (such as dislocations, stacking faults, microtwins) and therefore their ultra-high strength would explain the greater resistance to deformation than the amorphous phase itself. In metallic glasses, the deformation mechanism is known to be inhomogeneous (at room temperature) involving shear on highly localized bands [62,63]. Consequently, it has been suggested [57] that the very localized shear deformation can be effectively suppressed by the multiplication of interfaces (induced by the nanocrystallization) or by a possible interaction between shear bands and nanocrystallites (which would able to arrest their propagation). Thus, by the above considerations, multiplication of (strong) interfaces would be a key parameter as it is the case with conventional precipitation-hardened alloys. But in conventional fully crystalline alloys, hardening operates by hindering of dislocation motions, which can not be the case with the nanostructured composite materials presented here. Nevertheless, it can be supposed that during microhardness or abrasive tests, the transmission of the applied load (and the deformation) must be influenced by the structural division and separation at the nanoscale and have consequence on the strengthening effect. Nevertheless, in recent papers, Greer et al. [59,64,65] have shown that the hardness of fully amorphous or partially crystalline alloys correlates well with the solute content in the amorphous phase. From this observation, it is concluded that hardening mechanism of nanocrystallization is attributed more to the solute enrichment of the remaining amorphous matrix than to the presence of the nanoparticles themselves.
4. Conclusions Bulk Zr, Pd, La-based metallic glasses with a relatively low cooling rate can be processed by conventional casting while in order to produce Albased amorphous alloys, rapid solidification technique is still required. Metallic glasses can be devitrified by thermal treatment and nanostruc-
tured nanophase composites consisting of nanometer-sized crystallites (or quasicrystallites) embedded in an amorphous matrix are obtained. Fully amorphous metallic glasses are shown to behave similar as hardened-alloys by comparing them in a wear resistance/hardness diagram of different material classes and represent useful properties for coating application. Hardness and wear resistance are both observed to be enhanced by nanocrystallization, particularly in the case of Al-based alloys. The reasons are discussed as connected with ultrahigh strength nanoparticles, multiplication of interfaces or hardening effect due to a solute enrichment of the remaining amorphous matrix.
Acknowledgements This work has been conducted through an international collaboration and A.L. Greer is particularly thanked for his constant help and participation and for the fruitful discussion on the subject. The author acknowledges the contributions to this work from K. Hono, D.H. Ping, A. Aronin, M.Y. Myung and A. Inoue.
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Microhardness and abrasive wear resistance of metallic glasses and nanostructured composite materials Thierry Gloriant
*
Laboratoire de M etallurgie et Physico-Chimie des Mat eriaux, G.R.C.M. (EA2620) Institut National des Sciences Appliqu ees de Rennes 20, avenue des Buttes de Co€esmes, F-35043 Rennes cedex, France
Abstract The present paper reports on as-cast bulk zirconium-based, lanthanum-based and palladium-based metallic glasses and on melt-spun aluminum-based amorphous ribbons. Their glass forming abilities, nano(quasi)crystallization mechanisms and mechanical properties are discussed through a short review of very recent papers and by the presentation of some experimental results. The tribological aspect such as microhardness and abrasive wear resistance are more particularly investigated for their potential applications technologically useful as coatings. The formation of nanostructured composite materials consisting of nanophase particles embedded in a residual amorphous matrix by appropriate thermal treatments is also studied and the influence of the nanocrystallized volume fraction is evaluated and discussed for both hardness and wear resistance. Ó 2003 Elsevier Science B.V. All rights reserved. PACS: 81.05.K; 61.43.D; 62.25
1. Introduction Since the first success of preparing an amorphous phase in the Au–Si system by rapid solidification in 1960 [1] a great number of metallic amorphous alloys has been produced until now. Before 1990, Fe-, Co- and Ni-based amorphous alloys were formed by using high cooling rates but the resulting sample thickness was limited to less than about 50 lm [2] reducing their potential as engineering materials. However, in 1984, an ex-
*
Tel.: +33-2 23 23 85 19; fax: +33-2 23 23 82 40. E-mail address: [email protected] (T. Gloriant). URL: http://www.insa-rennes.fr/l-chmet/accueil.htm.
ceptional example, Pd40 Ni40 P20 was discovered as the first bulk metallic glass (BMG) forming alloy with a critical cooling rate as low as about 1 K s1 when processed in B2 O3 [3]. Thus for a long period before 1990, no other amorphous alloys except the Pd–Ni–P and Pt–Ni–P systems [4] were synthesized because of the necessity of high cooling rate above 105 K s1 for glass formation. The discovery of the La- [5] and Zr-based metallic glasses [6,7] in the late 1990s from multicomponents systems has triggered tremendous research activity in this area particularly for their excellent mechanical properties. Thus, a number of bulk amorphous alloys have been produced in Zr–Al–TM [6,8], Zr–Ti– TM–Be [7], Ti–Zr–Al–TM–Be [9], Pd–Cu–Ni–P [10] or in La–Al–TM systems [5] (TM ¼ transition metal) by conventional casting processes. It is to
0022-3093/03/$ - see front matter Ó 2003 Elsevier Science B.V. All rights reserved. PII: S 0 0 2 2 - 3 0 9 3 ( 0 2 ) 0 1 9 4 1 - 5
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be noticed however that the maximum sample thickness reaches about 40 mm maximum for Zrbased alloy [11,12] and 70 mm for Pd-based alloy [13]. In order to produce metallic glasses from the melt, the cooling rate must be sufficiently high to avoid nucleation and growth of crystalline phases in the supercooled liquid region between melting temperature Tm and glass transition Tg . As recently reported by Schneider [14], Turnbull [15], using a simple nucleation theory, showed that the reduced glass transition temperature, Trg ¼ Tg =Tm , is a key parameter that determines whether or not the melt of a given material can form a glass during cooling and concluded that to produce bulk metallic glasses, Trg must be relatively high. On the other hand, in order to develop some of the best bulk glass-forming alloys, Johnson [16] has successfully added suitable components (Cu, Ni) to lower the eutectic temperature in the Zr–Ti–Be system obtaining the family of Zr–Ti–Cu–Ni–Be glasses. These alloys show an exceptional glass forming ability and a critical cooling rate as low as about 1 K s1 such as for Vitreloye, which makes them very interesting materials for engineering applications [7]. With the same objective and in order to optimize the stabilization of metallic supercooled liquid and bulk amorphous alloys, Inoue [17] has proposed the following three empirical rules to achieve bulk glass formation: (1) multicomponent systems consisting of more than three elements; (2) significant difference in atomic size ratios above about 12% among the three main constituent elements; and (3) negative heats of mixing among the three main constituent elements. By choosing the above-described multicomponent alloy systems, it is possible to produce amorphous ribbons by melt-spinning (when a very high cooling rate is required), and bulk amorphous alloys by solidification techniques such as copper mold casting, arc melting suction casting, squeeze casting, high pressure die casting, unidirectional melting or water quenching. The standard procedures are given in detail in [18]. Metallic glasses have also attracted increasing interest as precursors to produce nanophase composite materials because of the appearance of good properties such as mechanical [19,20] or
97
magnetic [21,22]. This was particularly investigated for Zr-based BMG and also for Al-based metallic glasses after their first elaboration by rapid solidification (melt spinning) in 1988 [23–25]. These amorphous alloys crystallize into a nanostructured composite material consisting of nanometer-sized crystallites (or quasicrystallites) embedded in an amorphous matrix, often further enhancing mechanical properties above those of the metallic glasses [19,20]. Some of the best alloys concerning the glass forming ability of Al-based amorphous alloys (>85 at.%Al) are those elaborated from Al–TM–RE (TM ¼ transistion metal, RE ¼ rare earth) systems [26,27]. However, the glass forming ability is not high enough to obtain them as BMG by conventional casting and unfortunately a rapid solidification process (>105 K s1 ) is still required up to now [28,29]. The devitrified nanostructure can be formed by the suppression of growth reaction caused by segregation of a solute element with low atomic diffusivity at the nanocrystal/amorphous interface and by the higher thermal stability of the remaining amorphous phase due to the solute enrichment occurring during crystallization. The final nanostructure is strongly dependent on the number of stable homogeneous nucleation sites in the as-quenched glass but it has also been suggested a possible control by a heterogeneous nucleation mechanism by additional elements occurring by formation of clusters or by the presence of quenched-in nuclei [30]. In Zr-based amorphous alloys, it has been reported that an icosahedral quasicrystalline phase is formed in some alloys with a supercooled liquid region and high glass forming ability [31,32]. Recently, K€ oster et al. [33] have reviewed the conditions for icosahedral quasicrystal formation upon annealing in Zr–Al–Ni–Cu bulk metallic glasses and the beneficial influence of alloying elements such as Hf or Y. In similar compositions, addition of Ag or Pd to the Zr–Al–Ni–Cu glassy alloys has been observed to be effective for the increase of the number of nucleation sites and for the suppression of grain growth of the nano icosahedral phase [34]. As previously mentioned, devitrification of Albased amorphous ribbons upon annealing leads to the formation of a nanometer-scale dispersion of
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a-Al crystallites (or nanoquasicrystallites, see [35] for review) in the remaining matrix and the nanocrystallization mechanism occurring by nucleation and growth was intensively studies by several authors [26,27,36]. Recently, it was observed by TEM that the crystallization is preceded by a phase separation of the amorphous matrix (of about 40 nm) and nucleation of a-Al nanocrystals starts preferentially at the interface of the phaseseparated regions [37]. Because of the limited thickness obtained in the preparation of metallic glasses, particularly with Al-based alloys, this paper is focused on the attractive results for tribological applications, which may suggest metallic glasses technologically useful as coating by a combination of high hardness and high abrasive wear resistance.
microscopy (3DAPFIM) in order to obtain chemical distribution maps of the elements (Al, Ni and Sm) at the nanometer scale (see [41] for more details concerning this study). The 3DAPFIM technique (described in Ref. [42]) has been proven to be ideal for studying solute clustering, precipitation, segregation or partitioning in nanocrystallization [42]. For the wear resistance measurements, a microscale abrasion test was used. This involves rotating a steel sphere against a test specimen in the presence of small angular SiC particles. This mode of abrasion is commonly termed Õthree-bodyÕ and is described in more details elsewhere [43–45]. The wear resistance may be calculated from the slopes of the curves representing the wear crater volume measured during the abrasion test following the classical model proposed by Rabinowitz et al. [46] according to
2. Experimental methods Vw ¼ This paper presents experimental results on three BMG (Zr-based, Pd-based and La-based) and two Al-based amorphous ribbons. Bulk amorphous Zr65 Al7:5 Ni10 Cu17:5 (16 mm of diameter) and Pd40 Ni10 Cu30 P20 (40 mm of diameter) compositions were prepared by water quenching using the standard procedure described in detail elsewhere [10,38] and the bulk La60 Al20 Ni10 Co5 Cu5 glass (2 mm of diameter) was prepared by a high-pressure die casting method as presented in [39]. Amorphous Al88 Ni4 Sm8 and Al86 Ni11 Yb3 alloys were produced by melt spinning and the ribbons obtained were typically 3 mm wide and 50 lm thick. The volume fraction of nanocrystalline a-Al transformed during different annealing times (in tubular furnace under argon atmosphere) was evaluated using the enthalpy of formation calculated from the first differential scanning calorimetry peak according to the method proposed in Ref. [40]. The as-quenched amorphous state and the crystallization behavior under annealing were characterized by X-ray diffraction and by transmission and high resolution electron microscopy (TEM and HREM) using a Philips CM30 microscope. The Al88 Ni4 Sm8 partially crystallized alloy was particularly investigated by three-dimensional atom-probe field-ion
SN ; Rw
ð1Þ
where Vw is the wear volume measured with a calibrated optical microscope, S, the sliding distance, N , the normal load and Rw , the wear resistance. Vickers testers were used with a 0.20 or 0.49 N load for the microhardness measurements.
3. Results and discussion 3.1. Hardness and wear resistance of metallic glasses It has been described in Section 1 that the glassy alloy exhibit high strength combined with high toughness leading to interesting practical use as engineering materials. For example, bulk La–Al– TM (TM ¼ transition metal) alloys exhibit high tensile strength [39], Zr-based amorphous alloys present superplastic deformation [47] and high impact fracture energy [48] and Al–Ni–RE (RE ¼ rare earth metal) amorphous ribbons exhibit high tensile strength above 980 MPa and good bending ductility [24]. For these reasons, metallic glasses have been extensively studied but there is not a lot of work concerning the wear resistance.
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There are many possibilities to evaluate the abrasive properties of materials such as pin-on disk, pin-on plate, cylinder, ball or wheel on flat or scratch test. . . The first documented wear test on a metallic glass was by Boswell in 1997 who studied the abrasive wear of Pd78:1 Si61:4 Cu5:5 by using a pin-on-disk set-up [49] but most of the studies in this field were concentrated on iron-based alloys because these were the first to become available in industrial quantities and to present a relatively good wear resistance [50]. By using the technique presented in this paper, Greer et al. were the first to present results concerning Al, Zr, Pd and Labased amorphous alloys [51–53]. Fig. 1 displays the volume of the wear craters (measured as a function of the sliding distance) in metallic glasses representing four common types of glass compositions, i.e. an Al-based amorphous ribbon and bulk amorphous zirconium, lanthanum and palladium based alloys as described in Section 2. The measurements were carried out by interrupted tests as a function of sliding distance and show that Eq. (1) is obeyed by each of the metallic glasses. Fig. 2 compares the wear resistance of the four metallic glasses in correlation with their microhardness values and compared with a wide variety of materials by using the same micro-scale abrasive wear test. It can be seen that the wear resistance in-
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Fig. 2. The correlation of abrasive wear resistance with microhardness for different material classes (measured using the same technique as in the present work). The four glasses are represented by open squares [44,45].
creases linearly with the hardness, as expected within a given material class. This comparison of materials, as observed in Fig. 2, indicates that the metallic glasses fit firmly in the classification of hardened alloys. For a given hardness, hardened alloys show wear resistances significantly greater than those of ceramics (which are reduced by brittle failure mechanisms) and significantly less than those of pure metals (which are enhanced by work hardening). The wear resistance of metallic glasses, much greater than that of the ceramics of the same hardness, suggests that the wear mechanism does not involve brittle fracture, which is of great interest for coating applications. 3.2. Nanocrystallization of metallic glasses
Fig. 1. The volume wear craters in the four metallic glasses measured as a function of sliding distance during interrupted tests.
An example of a devitrified nanostructure is shown in the TEM dark-field image presented in Fig. 3. The micrograph shows the dispersion of aAl nanocrystallites (in bright) of about 20–30 nm in size precipitated after a thermal treatment at 500 K during 120 min of the Al88 Ni4 Sm8 amorphous composition. This composition was also studied by high-resolution electron microscopy (see the HREM micrograph in Fig. 4 showing a nanocrystal/amorphous interface) and by threedimensional atom-probe field-ion microscopy
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partitioning is confirmed to play a significant role in impeding their growth and aiding the formation of a fine-scale microstructure by the rejection of the two solute elements from the pure a-Al crystals to the matrix during crystallization. 3.3. Effect of the nanocrystallization on hardness and wear resistance
Fig. 3. TEM dark-field image of Al88 Ni4 Sm8 amorphous alloy heat-treated at 500 K for 120 min.
Fig. 4. High-resolution transmission electron microscope image (HREM) of Al88 Ni4 Sm8 (annealed: 500 K/120 min).
(3DAPFIM) (results have been reported in a recent paper [41]). As presented in Fig. 5, it is observed by 3DAPFIM that nanocrystallites appear to be completely devoid of solute content (Ni and Sm). The three-dimensional map on the left side of Fig. 5 represents the samarium content in the melt (each point is one Sm atom) and shows a region devoid of Sm atom. This zone corresponds to the 100 at.% aluminium isoconcentration surface presented on the right 3D map in the figure showing one pure a-Al nanocrystal. Consequently, solute
It is well established that mechanical properties are further enhanced by homogeneous dispersion of nanoscale compound particles into the glassy phase. For instance, Inoue et al. [54] have observed that precipitation of nanoscale quasicrystalline particles with a size of about 30 nm in the Zr–Al– Ni–Cu–(Ag or Pd) bulk glassy alloy exhibits enhanced mechanical properties such as higher tensile or compressive strength, larger elongation an higher hardness and Eckert et al. [55] have observed high strength and good ductility in nanostructured Zr-based multiphase alloys. For a La-based BMG, recent experimental results show that introducing nanocrystalline particles by way of heat treating the amorphous alloy can increase its stiffness and strength [56]. On the other hand, Al-based nanophase composites are known to present quite remarkable tensile strength at room temperature as high as about 1.5 times that of the corresponding fully amorphous alloy and about 3 times that of a conventional precipitation-hardened alloy [57]. For tribological interest, it was shown that hardness and wear resistance are also enhanced by nanometer-scale dispersion of crystallites or quasicrystallites in an amorphous matrix. By studying nano-scratch behavior of bulk Zr-based alloys, Wang et al. [58] have recently shown that samples having a structure consisting of an amorphous matrix with dispersed nanocrystalline particles exhibits the best wear performance. In amorphous alloys from Al–Ni–RE systems, hardness and abrasive wear resistance were also enhanced by nanocrystallization [51,59– 61]. Fig. 6 presents experimental results obtained for the melt spun Al86 Ni11 Yb3 amorphous ribbon. This figure reveals a clear correlation between the microhardness and the volume fraction of a-Al nanocrystallites obtained by isothermal annealing at 473 K. The microhardness is shown to increase
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Fig. 5. Three-dimensional Sm atom distribution (by 3DAPFIM) after annealing Al88 Ni4 Sm8 at 500 K for 15 min (left). Corresponding 100 at.% Al isoconcentration surface (right) [28].
Fig. 6. Volume fraction of a-Al nanoparticles crystallized and microhardness of Al86 Ni11 Yb3 alloy for different times of annealing at 473 K [61].
more than twice, reaching about 400 Hv for a crystallized volume fraction of about 25%. The wear resistance is also observed to be enhanced with the a-Al volume fraction crystallized as shown the Fig. 7. In this figure are presented wear resistance and volume fraction curves of the Al88 Ni4 Sm8 alloy isothermally annealed at 500 K.
Fig. 7. Volume fraction of a-Al nanoparticles crystallized and abrasive wear resistance of Al88 Ni4 Sm8 alloy for different times of annealing at 500 K [60].
After about 120 min of heat treatment, the wear resistance has increased from 1.5 to 2.3.10þ11 Pa induced by a nanocrystalline volume fraction of about 40%. The mechanism that can explained the observed increase of mechanical properties such as hardness and wear resistance by formation of nanophase
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particles in a residual amorphous matrix has been extensively discussed. Kim et al. [57] have pointed out that the nanocrystallites are too small to content defects (such as dislocations, stacking faults, microtwins) and therefore their ultra-high strength would explain the greater resistance to deformation than the amorphous phase itself. In metallic glasses, the deformation mechanism is known to be inhomogeneous (at room temperature) involving shear on highly localized bands [62,63]. Consequently, it has been suggested [57] that the very localized shear deformation can be effectively suppressed by the multiplication of interfaces (induced by the nanocrystallization) or by a possible interaction between shear bands and nanocrystallites (which would able to arrest their propagation). Thus, by the above considerations, multiplication of (strong) interfaces would be a key parameter as it is the case with conventional precipitation-hardened alloys. But in conventional fully crystalline alloys, hardening operates by hindering of dislocation motions, which can not be the case with the nanostructured composite materials presented here. Nevertheless, it can be supposed that during microhardness or abrasive tests, the transmission of the applied load (and the deformation) must be influenced by the structural division and separation at the nanoscale and have consequence on the strengthening effect. Nevertheless, in recent papers, Greer et al. [59,64,65] have shown that the hardness of fully amorphous or partially crystalline alloys correlates well with the solute content in the amorphous phase. From this observation, it is concluded that hardening mechanism of nanocrystallization is attributed more to the solute enrichment of the remaining amorphous matrix than to the presence of the nanoparticles themselves.
4. Conclusions Bulk Zr, Pd, La-based metallic glasses with a relatively low cooling rate can be processed by conventional casting while in order to produce Albased amorphous alloys, rapid solidification technique is still required. Metallic glasses can be devitrified by thermal treatment and nanostruc-
tured nanophase composites consisting of nanometer-sized crystallites (or quasicrystallites) embedded in an amorphous matrix are obtained. Fully amorphous metallic glasses are shown to behave similar as hardened-alloys by comparing them in a wear resistance/hardness diagram of different material classes and represent useful properties for coating application. Hardness and wear resistance are both observed to be enhanced by nanocrystallization, particularly in the case of Al-based alloys. The reasons are discussed as connected with ultrahigh strength nanoparticles, multiplication of interfaces or hardening effect due to a solute enrichment of the remaining amorphous matrix.
Acknowledgements This work has been conducted through an international collaboration and A.L. Greer is particularly thanked for his constant help and participation and for the fruitful discussion on the subject. The author acknowledges the contributions to this work from K. Hono, D.H. Ping, A. Aronin, M.Y. Myung and A. Inoue.
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