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Effect of ball milling parameters on the microstructure of W–Y powders and sintered samples
International Journal of Refractory Metals & Hard Materials 21 (2003) 205–213 www.elsevier.com/locate/ijrmhm
Effect of ball milling parameters on the microstructure of W–Y powders and sintered samples M.-N. Avettand-Feno€el *, R. Taillard, J. Dhers 1, J. Foct Laboratoire de M etallurgie Physique et G enie des Mat eriaux, UMR CNRS 8517, Universit e de Lille, F-59655 Villeneuve d’Ascq, France Received 24 March 2003; accepted 11 June 2003
Abstract High energy milling carried out in a planetary ball mill was used in order to alloy elemental powders and to obtain nanostructures and oxide dispersion strengthened (ODS) alloys. Owing to such advantages, this process was retained so as to elaborate ODS tungsten alloys from tungsten and yttrium powders by using a WC–Co milling system with 16 balls. The experiments were performed for a duration of up to three days while applying a 400 rpm speed and a ball-to-powder weight ratio of 16. The W–1 vol%Y blends were subsequently compacted at room temperature and sintered at 1800 °C for 4 h. The present paper deals with the effect of milling time on the mechanical behavior of the powders, on the refinement of the microstructure, on the improvement of the second phases distribution and on the contamination by cobalt, carbon and oxygen. At last, these results are related with the behavior of the powders during densification. Ó 2003 Elsevier Ltd. All rights reserved. Keywords: Tungsten alloy; ODS alloy; Ball milling; Sintering; Microstructure
1. Introduction The aim of the study is to elaborate novel materials with an enhanced ductility at high temperature. An oxide dispersion strengthened (ODS) tungsten alloy prepared by sintering of reactively ball milled blends of elemental tungsten and yttrium was retained in order to fulfil such a goal. Such a choice reposes on the following background: tungsten is the most refractory metal. It is nevertheless embrittled by the grain boundary segregation of low levels of oxygen [1]. Now, according to many works, ODS alloys are the well suited candidates to resist to high temperature creep [2–4]. Therefore the idea consists in producing ODS tungsten base alloys, while taking advantage of the high chemical affinity of yttrium with oxygen. Such a reactivity actually overwhelms that of tungsten with oxygen [5,6]. Reduced amounts of pure yttrium powder are then mixed with an elemental tungsten powder in order to trap the oxygen contained in tungsten during ball milling. Another reason to select yttrium is that W–Y2 O3 alloys present a good corrosion resistance [7]. Powder metallurgy methods and more * 1
Corresponding author. CEA Valrh^ o, B.P. 111, F-26702 Pierrelatte cedex, France.
0263-4368/$ - see front matter Ó 2003 Elsevier Ltd. All rights reserved. doi:10.1016/S0263-4368(03)00034-9
particularly ball milling and sintering were indeed used so as to elaborate the alloys. Ball milling gives rise to fine and well dispersed oxides in the structure. Many ODS superalloys have actually been produced using this process [2–4,8]. In addition, ball milling takes the advantage of generating nanostructures, which make the sintering easier [9]. First experiments on tungsten and tungsten– yttria powder blends [10,11] have already been achieved in order to define suitable parameters. The present paper depicts the effects of milling time on the microstructure and properties of the powders and sintered alloys primitively composed of pure tungsten and yttrium.
2. Experimental Tungsten powder (190 ppm O, 15 ppm C and 5 lm average particle size) and yttrium particles (10 ppm maximum Fe2 O3 and maximum particle size of 250 lm) were mixed with a nominal composition of W–1 vol%Y in a Fritsch Pulverisette 6e planetary ball mill. Ball milling was carried out for different times up to three days in a WC–6 wt% Co system composed of a vial and of 16 balls of 10 mm diameter. A ball-to-powder weight ratio of 16:1 and a 400 rpm speed were adopted. The loading
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and unloading of the powders into and from the containers were always done inside an argon filled glove box. The powder blends were axially compacted at room temperature and under air with a pressure of 1.2 GPa. The pellets were subsequently sintered under vacuum (<10 2 Pa) by using a 300 °C/h heating rate followed by a 4 h heat treatment at 1800 °C completed by a 500 °C/h cooling stage. Laser diffractometry experiments were conducted on the powders to determine their size. Chemical analyses of the powders was performed, on the one hand by an inductively coupled plasma (ICP) atomic emission spectrometer to measure the cobalt content, and on the other hand by burn up using ELTRA ON-900e and LECO CS244e devices so as to evaluate carbon and oxygen contents. The microstructure, i.e. the particle morphology, the nature and the distribution of the second phases, the contamination by the milling system, was more especially studied by means of field emission gun/scanning electron microscopy (FEG/SEM), X-ray diffractometry (XRD) with a cobalt anticathode, electron probe microanalysis (EPMA) and transmission electron microscopy (TEM/STEM) analyses. The lattice strain and crystallites size were estimated from XRD results, applying the Warren–Averbach method [12] and using the {1 1 0} tungsten Bragg peaks. In the present paper, the word ÔcrystalliteÕ is going to designate a coherent diffraction domain within a single particle of powder whereas the term ÔgrainÕ is reserved for the sintered samples, where it keeps its usual meaning within bulk samples. The densities of the bulk sintered samples were measured by the ArchimedeÕs water immersion method while their Vickers microhardness was obtained under a 100 g load.
3. Results 3.1. Microstructure 3.1.1. Morphology and size of the powder particles Fig. 1a–c are scanning electron micrographs of the asreceived W and the W–1 vol%Y powder milled for 20 h,
Fig. 2. Laser diffractometry analysis of the particle size distribution of the W–1%Y powders milled for different times.
respectively. The third micrograph is typical of the case of a long-milled powder. The elemental tungsten powder is characterized by clusters of polyhedral particles with an average size of 4.5 lm (Fig. 1a). Such a morphology is consistent with tungsten powder arising from the reduction of WO3 by hydrogen [13]. Fig. 1c proves that ball milling has significantly changed the shape of the powder particles which becomes more equiaxed at long milling times. At the same time, the particle size is significantly reduced. Fig. 2 further shows that after 5 min of milling, welding seems to predominate whereas at the longest milling times (20 h), the mean particle size is finer and becomes submicrometric. It can be guessed, from such an observation, that fracture is the major mechanism at long milling times. This evolution of particle size is confirmed by bidimensional measurements of the powder particles on scanning electron micrographs [14,15]. Furthermore, it should be underlined that the milling time does not exert the same effect on the evolution of the size of the particles and of the crystallites (Section 3.1.2). 3.1.2. Inner microstructure of the powder particles and of the sintered samples Fig. 3a and b indicate the effect of milling time on the crystallite size and on the lattice strain of the samples measured along the [1 1 0] direction, respectively.
Fig. 1. Morphology of the elemental tungsten powder (a and b) and of the W–1%Y blend milled for 20 h, (c) BSE/FEG/SEM.
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Fig. 3. Effect of milling time on crystallite size (a) or on lattice strain (b) (XRD results).
For the powders, the longer the milling time, the finer the crystallite size is, up to a constant value. The present evolution suggests that the steady state of the process of ball milling is certainly attained from about 30 h of milling. Moreover it is worthy to note that, in this condition, the mean crystallite size is close to 5 nm. After sintering and at a given milling time, the full width at half maximum of the (1 1 0) peaks of tungsten diminishes, which indicates a growth of the grains during sintering (Fig. 3a). However the XRD results markedly differ from the measurements supplied from the FEG/SEM micrographs. For instance, Fig. 8c displays that the average grain size of an initially milled for 10 h sample grows to 10 lm after sintering, which is quite different from the value of 70 nm indicated by XRD results. Such a difference could be linked with two facts. Firstly, the XRD measurements are actually made in a preferential direction. Such a remark seems however to be of second importance because of the equiaxiality of the grains in the sintered samples. A second and more likely explanation lies in the fact that MurakamiÕs reagent etching only exhibits highly misoriented grain boundaries. Further works are necessary so as to be more conclusive on this topic. Contrary to the crystallite size, the lattice strain of the powder particles globally rises with the milling time to finally tend towards 0.3% (Fig. 3b). Ball milling generates many defects in the powder particles, which explains the growth of the average lattice strain. Nevertheless and in accordance with the literature data, the value of this lattice strain is always rather low and under its 0.58% estimated for pure milled tungsten [16]. The comparison with this upper result is not easy owing to different milling conditions.
phase on the X-ray diffractograms seems to arise more from both the too low nominal yttrium content and the marked absorption of the yttrium or yttria X-rays reflections by tungsten than from the fineness of the yttrium phase. Indeed, for the shortest milling times, the distribution of yttrium rich particles is rather coarse i.e. their spacing and their size are of about 7 and 3 lm, respectively. Moreover, according to the observations made with back scattered electrons (BSE)/FEG/SEM, these yttrium rich particles are only visible after the milling times shorter than 80 min. At least, in the latter conditions, they behave as less ductile particles than the tungsten ones which are flattened against their surface (Fig. 4). In the same way and as shown by Fig. 5, the EPMA results display a discrete distribution of yttrium rich phase within the particles of the W–1 vol%Y blend milled for 80 min. Moreover, it is significant to note that the same figure shows that such particles are not always oxidized. Such an observation also agrees with the lower than 624 ppm oxygen content of the powders milled for less than 40 min (Fig. 6c). The latter value is required in order to entirely change yttrium into yttria. The bulk metallic or Y(O) solid solution nature of this not oxidized yttrium can however be questioned. Otherwise, according to Fig. 6c, it is possible that yttrium is
3.1.3. Second phases This paragraph deals with the nature and the dispersion of the various kinds of second phases, including the yttrium rich ones. 3.1.3.1. Case of the blended powders. For any milling time, the absence of detection of an yttrium rich second
Fig. 4. W–17%Y powder milled for 80 min (BSE/FEG/SEM). Some white tungsten flakes can be seen over the surface of the grey yttrium rich particle.
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Fig. 5. Weight contents of oxygen, yttrium and tungsten within a large powder particle of W–1%Y milled for 80 min: (a) secondary electron image (the arrow indicates the position of the line of analysis) and (b) profiles of chemical composition.
Fig. 6. Effect of milling time on the contamination in mass percent of the powders by carbon (a), cobalt (b) and oxygen (c).
completely oxidized at the longest milling times. By way of contrast, the EPMA study failed to detect an yttrium rich phase in the W–1 vol%Y powders milled for 30 h. Such a result suggests that the yttrium rich phase is very fine in the final stationary stage of ball milling. The contamination of the blends has further been quantified. The chemical analyses show that the carbon and the cobalt contents of the blends continuously increase with the milling time (Fig. 6a and b). Moreover, and a priori in accordance with this observation because tungsten carbide and cobalt are the constituents of the milling system, Fig. 7a depicts a marked increase of the volume fraction of WC in the same powder for longer than 4 h milling times. These estimations were deduced from the measurements of the areas of the WC X-ray diffraction peaks (Fig. 7b) supposing both the complete transformation of yttrium into yttria and a value of the WC/Co ratio that is the same as the composition of the milling media. 3.1.3.2. Case of the sintered samples. Fig. 8a and b show that the longer the milling time, the more homogeneous and finer the second phases distribution. For any milling time, the size of the yttrium rich particles is however rather variable. For instance, and as depicted in Figs. 8c and 9, it is comprised between a few tens of nanometers and a few micrometers after the sintering of the blends milled for 20 h. Fig. 8c further shows that the second phases are essentially located at the grain boundaries. The EPMA analyses of the W–1 vol%Y sample sintered from the milled for 30 h powder
Fig. 7. (a) Contamination of the W–1%Y powders by the WC/Co milling system (XRD) and (b) example of X-ray diffraction pattern of the W–1%Y powder blend milled for 20 h.
mixture (Fig. 10) indicate that these second phases possess a stoichiometry close to yttria. The results are checked by the electron diffraction patterns of some nanometric white dispersoids like those shown in Fig. 9.
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Fig. 8. Dispersion of second phases in the W–1%Y sintered sample obtained from the blend milled for 40 min (a), 20 h (b) or 10 h (c). The latter sample was previously etched by the Murakami reagent in the latest case (BSE/FEG/SEM).
Fig. 9. STEM dark field of the yttria dispersoids in the W–1%Y sample sintered from the blend milled for 20 h.
Their crystal lattice is actually suited to such a chemical composition. Concerning the contamination of the sample sintered from the powder milled for 30 h, the STEM/EDX analyses indicate that the cobalt content is equal to 1.5 wt% in average in the matrix while it can attain values of 9.4 wt% in cobalt rich grains. 3.1.4. Properties of densified samples Fig. 11 shows that the green density is reduced from about 70% to 51% with the milling time. Moreover, up to 80 min of milling, the green density and the sintered density decrease in the same way while,
Fig. 10. EPMA analysis of the nature of an yttrium rich second phase in the W–1%Y sintered sample made of the milled for 30 h powder. The arrow indicates the position of the line of analysis.
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Fig. 11. Effect of milling time on both green and sintered density.
at longer milling times, the sintered density is improved contrary to the green density.
4. Discussion 4.1. Mechanism of ball milling In case of a ductile/ductile behavior of powders, the particles are at first flattened. This early stage is followed by a period dominated by particle welding which gives rise to lamellar particles. Afterwards the particles become more equiaxed, with first an oriented lamellar morphology, and subsequently a random arrangement of several lamellar granules within a given particle. Finally, a steady state is reached. It is characterized by a refinement of the inner microstructure while the size distribution of the particles remains constant [17,18]. The present evolution of the size and of the morphology of the particles agrees with Figs. 1 and 2 while it is more detailed in another paper [15] that clearly establishes that the W–1 vol%Y blend behaves as a ductile/ductile system during milling. 4.2. Powder microstructure The XRD 5 nm size of the crystallites inside the powder particles milled for the longest durations agrees with the literature data. At the steady state, a mean size of 10 nm has indeed been measured for pure tungsten [16]. Such an agreement is observed in spite of changes in the milling conditions and in the initial size and purity of the powder mixture. Moreover, such values correspond to the 4.5, 5 [16,19] and 9 nm [20] theoretical estimations of dislocation free crystallites in pure tungsten. Concerning the contamination, the carbon content accounts for both the carbon adsorption over the powder surfaces and the contamination of the powders by the milling system which is made of WC/Co. Figs. 6
and 7 prove that the latter contamination is very important. It very likely arises from an adhesive wear of the milling media because of both the tendency of cobalt and tungsten carbide to adhere [21] and the overall ductile behavior of the powder blends during milling [22]. Moreover, the values of cobalt and carbon obtained by chemical analyses agree with the results of X-ray diffractometry concerning the tungsten carbide content (Figs. 6 and 7). It proves that the carbon content in the powders is essentially supplied by the tungsten carbide of the milling system. The cobalt content measured by chemical analyses in the W–1 vol%Y sintered sample agrees with the decrease of the tungsten lattice parameter measured by XRD in the original blend to which decreases from 3.1648 A 3.1559 A after 30 h of milling. Such a diminution agrees with its estimate by using both the VegardÕs law and the chemical determination by burning up of the cobalt content in the powder. According to Fig. 6, the oxygen content increases with the milling time to attain very large values paying regard to the argon composition of the milling atmosphere. The origin of this oxygen and of its large content must therefore be questioned and the more as oxygen embrittles the tungsten [1]. The reactivity of the powders towards the traces of oxygen present within the milling atmosphere is very likely promoted by the fineness of the powder particles. However, the vial is not perfectly sealed as it has been checked by changing the chemical composition of the milling atmosphere that oxygen arises essentially from a leak. Likewise it is interesting to note that the oxygen content increases with the milling time. Many second phase particles have an effect on the plasticity of the blend. Firstly, the yttrium rich phases seem less ductile than the tungsten powder at least at early milling stages (Fig. 4), which is contrary to the literature data which mention a far higher ductility for the yttrium metal than for the tungsten one [23]. Such an anomaly is very likely to arise from the chemical composition of the yttrium rich particles, i.e. are they metallic or oxidized? According to both the present results and previous ones [15], a prolongation of the milling time promotes the fracture of tungsten particles, and increases the contamination by cobalt [which is metallic or oxidized (XPS experiments)] and tungsten carbide (Figs. 6b and 7a). Moreover, it can be suggested that for the longest durations of milling, yttrium is changed into yttria. The fracture of the tungsten particles could be explained by an enhancement of their plastic deformation and therefore of their work hardening that entails their fracture. Plastic deformation would then be favored by the stress concentration due to the hard particles of yttria and of tungsten carbide. According to its high volume fraction (Fig. 7), tungsten carbide plays a dominant
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role on the general behavior of the powder. In addition, the formation of a cobalt solid solution in the tungsten is likely to modify the mechanical behavior of the blend. 4.3. Densified samples The decrease of the green density with the milling time can be due to many factors. The green density actually depends on a great number of parameters dealing with the morphology, the size distribution and the ductility of the powder particles. Such a ductility is governed by both the distribution of the second phases within the compacts and the degree of contamination and the dislocation substructure of the powders. According to the present results concerning the density of the green compacts and the properties of the W–1 vol%Y milled powders, two assumptions may be considered. Firstly, the refinement of the particles size with the milling time gives rise to a narrow monomodal size distribution (Fig. 2) which is less favorable than a bimodal size distribution to eliminate the porosity within both a green compact [16] and a sintered sample [24]. Secondly, the rise of the contamination, for instance by oxygen (Fig. 6c), can impede the particles to weld during the stage of compaction. The oxide layer on the powder particles surfaces prevents them to deform and therefore contributes to a low densification. It is however not possible to check the latter hypothesis because of the easiness of the green compacts to crumble. Over 4 h of milling, the sintered density rises whereas the green density continues to decrease (Fig. 11). The origin of such a discrepancy is very likely to have to do with the nature of the sintering mechanism that can however be questioned. Some clues suggest that the sintering proceeds in solid phase whereas others could indicate that it is promoted by the formation of a liquid phase. In the following, both cases are successively considered. Solid phase sintering seems to occur because of the presence of polyhedral grains (Fig. 8b and c). Solid phase sintering is especially plausible because the sintering temperature is rather low for tungsten, which is usually solid phase sintered at temperatures around 2000 °C [24]. Moreover, the contamination by the milling system, i.e. by cobalt and tungsten carbide, is obvious from 4 h of milling, which suggests that cobalt has to do with the rise of the sintered density that occurs in spite of the decrease of the compactÕs one. The literature data actually indicate that the sintering of tungsten can be activated by the presence of cobalt in solid solution because this chemical element entails an increase of the diffusion rate [25–27]. Moreover, cobalt is actually in solid solution in the sintered sample insofar as the 1.5 wt% content of cobalt found by STEM/EDS in the W–1 vol%Y milled for 30 h and sintered sample agrees with the VegardÕs law.
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Alternatively, various features suggest that the sintering of tungsten can locally proceed in liquid phase. Firstly, the liquid phase may arise from the melting of isolated islands of the WC/Co cermet which is usually liquid phase sintered at about 1450 °C [27,28]. Secondly, pure yttrium and a 0–0.304 at% O–Y solid solution melt at 1522 °C [29] and 1522–1670 °C [30], respectively, that is to say at temperatures that are significantly lower than the 1800 °C present sintering temperature. The latter consideration agrees with both the presence of some rounded interfaces with the second phases (Fig. 8a) as well as with the growth of intergranular second phases during sintering. Such a growth has been noticed by the present authors comparing the microstructures of sintered samples of W plus 1 vol% or 17 vol%Y [15]. Itoh et al. [31] have also noticed the growth of the yttria particles and its increase with both the yttria content and the sintering temperature in W– Y2 O3 samples. Due to the very low solubility of yttrium in pure tungsten [29], such phenomena are rather astonishing and the more astonishing as, in the present case, the presence of cobalt cannot improve the solubility because of the tendency of yttrium and cobalt to form intermetallic compounds [32]. Thirdly, liquid phase sintering can further arise from the formation either of an 0.0785 at% W–Y eutectic at 1522 °C [31] or a W–Y2 O3 one at 1560 °C [30]. At last, liquid phase sintering may again proceed because of the formation of a Y2 (WO4 )3 phase which melts at 1440 °C. Such a compound has indeed been mentioned by Itoh et al. in W–Y2 O3 blends [31]. Consequently, numerous sintering mechanisms are very likely to intervene in a more or less exclusive way. Some further experiments would be necessary in order to be more conclusive. Otherwise, the formation of finer and finer nanostructures with the milling time usually improves the densification [9]. The comparison between Figs. 3 and 11 however indicates that this parameter is not predominant owing to a successive decline and rise of the sintered density with the prolongation of milling. At last, Fig. 3a, one more time, confirms that a significant grain growth is usual during the sintering of nanostructures (Section 3.1.2). Such a growth of the crystallites during sintering is however the less important, the longer the milling time (Fig. 3a) i.e. as the sizes of the particles and of the crystallites inside these particles decrease with the milling time. The whole of these observations suggest that the content and the distribution of the second phases are not suitable in order to fully stabilize a nanometric crystallite size of the powders by the Zener pinning mechanisms of grain boundaries [31,33] (Sections 3.1.3 and 3.1.4). However, after the longest milling times, the quality of the second phases distribution is more convenient than after the first steps of milling.
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4.4. Prospects The use of a WC/Co milling system brings about a high contamination of the samples by tungsten carbide and cobalt for the longest milling times, which could be avoided by changing the nature of the milling system. Such a contamination nevertheless presents an advantage as cobalt may improve the density of the sintered alloys (Section 4.3). Likewise an eventual change of the nature of the milling system could be questioned in order to maintain the stability of the nanostructure at the longest milling time because the sole yttrium content could then be insufficient to limit the grain growth. Moreover, it is worthy to control the milling atmosphere so as to avoid or at least to limit the contamination by oxygen and then to satisfy the proposed goal.
5. Summary and conclusions 1. The present study prominently establishes the interest of long milling times in order to elaborate ODS alloys and nanostructures, which is confirmed by many authors [34]. A prolongation of the milling duration generates firstly a refinement of the size of both the powder particles and of their crystallites and secondly an improved distribution of the yttrium rich phase (and of the second phases in general). At the longest milling time, it can be guessed that the second phases are suitably dispersed so as to impede the displacement of grain boundaries: the increase of the crystallite size measured by XRD actually remains small after sintering. 2. The use of a WC/Co milling system brings about a high contamination of the samples by tungsten carbide and cobalt for the longest milling times. 3. Otherwise, the high chemical affinity of the powders towards oxygen remains to be solved. 4. The microstructural analyses of the sintered samples can validate a solid phase sintering mechanism with local liquid phase sintering. The occurrence of liquid phase sintering is probably the cause of the opposite evolutions of the green compact and the sintered densities.
Acknowledgements The authors are very grateful to N. Llorca (CEA Saclay, France) for giving them the opportunity of using a gloves box, to C. Boyaval (IEMN, Lille, France) for the FEG/SEM micrographs, to J.F. Lartigue (CERMEP Grenoble, France) for the chemical analyses of the powders, to O. Duplessis (CEA Saclay, France) for the laser diffractometry analyses, to N. Lochet (CEA Saclay, France) for the sintering, to C. Fucili (CEA Valrh^ o,
France) for the EPMA analyses and to L. Gengembre (University of Lille, Laboratory of Catalysis, France) for the XPS analyses.
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[26] German RM. The effect of the binder phase melting temperature on enhanced sintering. Metall Trans A 1986;17A:903–6. [27] German RM. Sintering theory and practice. New York: John Wiley and Sons; 1996. [28] Wang Y, Delanoe A, Lay S, Pauty E, Allibert CH. Morphology and growth of WC grains in WC–Co Cermets. Materiaux 2002, CM 01-083. [29] Pandian S, Nagender Naidu SV, Rama Rao P. W–Y (tungsten– yttrium). Binary Alloy Phase Diagram. Metals Park, OH: ASM; 1990. p. 3531–2. [30] Carlson ON. The O–Y (oxygen–yttrium) system. Bull Alloy Phase Diagrams 1990;11(1):61–6. [31] Itoh Y, Ishiwata Y. Strength properties of yttrium-oxide-dispersed tungsten alloy. JSME Int J, Series A 1996;39(3):429–34. [32] Brandes EA, Brook GB, editors. Smithells metals reference book. 7th ed. Oxford: Heinemann–Butterworths; 1992. p. 11-208. [33] Chung KH, Lee J, Rodriguez R, Lavernia EJ. Grain growth behavior of cryomilled INCONEL 625 powder during isothermal heat treatment. Metall Mater Trans A 2002;33A:125–34. [34] Benjamin JS. Mech Alloying. Sci Am 1976;234(5):40–8.
Effect of ball milling parameters on the microstructure of W–Y powders and sintered samples M.-N. Avettand-Feno€el *, R. Taillard, J. Dhers 1, J. Foct Laboratoire de M etallurgie Physique et G enie des Mat eriaux, UMR CNRS 8517, Universit e de Lille, F-59655 Villeneuve d’Ascq, France Received 24 March 2003; accepted 11 June 2003
Abstract High energy milling carried out in a planetary ball mill was used in order to alloy elemental powders and to obtain nanostructures and oxide dispersion strengthened (ODS) alloys. Owing to such advantages, this process was retained so as to elaborate ODS tungsten alloys from tungsten and yttrium powders by using a WC–Co milling system with 16 balls. The experiments were performed for a duration of up to three days while applying a 400 rpm speed and a ball-to-powder weight ratio of 16. The W–1 vol%Y blends were subsequently compacted at room temperature and sintered at 1800 °C for 4 h. The present paper deals with the effect of milling time on the mechanical behavior of the powders, on the refinement of the microstructure, on the improvement of the second phases distribution and on the contamination by cobalt, carbon and oxygen. At last, these results are related with the behavior of the powders during densification. Ó 2003 Elsevier Ltd. All rights reserved. Keywords: Tungsten alloy; ODS alloy; Ball milling; Sintering; Microstructure
1. Introduction The aim of the study is to elaborate novel materials with an enhanced ductility at high temperature. An oxide dispersion strengthened (ODS) tungsten alloy prepared by sintering of reactively ball milled blends of elemental tungsten and yttrium was retained in order to fulfil such a goal. Such a choice reposes on the following background: tungsten is the most refractory metal. It is nevertheless embrittled by the grain boundary segregation of low levels of oxygen [1]. Now, according to many works, ODS alloys are the well suited candidates to resist to high temperature creep [2–4]. Therefore the idea consists in producing ODS tungsten base alloys, while taking advantage of the high chemical affinity of yttrium with oxygen. Such a reactivity actually overwhelms that of tungsten with oxygen [5,6]. Reduced amounts of pure yttrium powder are then mixed with an elemental tungsten powder in order to trap the oxygen contained in tungsten during ball milling. Another reason to select yttrium is that W–Y2 O3 alloys present a good corrosion resistance [7]. Powder metallurgy methods and more * 1
Corresponding author. CEA Valrh^ o, B.P. 111, F-26702 Pierrelatte cedex, France.
0263-4368/$ - see front matter Ó 2003 Elsevier Ltd. All rights reserved. doi:10.1016/S0263-4368(03)00034-9
particularly ball milling and sintering were indeed used so as to elaborate the alloys. Ball milling gives rise to fine and well dispersed oxides in the structure. Many ODS superalloys have actually been produced using this process [2–4,8]. In addition, ball milling takes the advantage of generating nanostructures, which make the sintering easier [9]. First experiments on tungsten and tungsten– yttria powder blends [10,11] have already been achieved in order to define suitable parameters. The present paper depicts the effects of milling time on the microstructure and properties of the powders and sintered alloys primitively composed of pure tungsten and yttrium.
2. Experimental Tungsten powder (190 ppm O, 15 ppm C and 5 lm average particle size) and yttrium particles (10 ppm maximum Fe2 O3 and maximum particle size of 250 lm) were mixed with a nominal composition of W–1 vol%Y in a Fritsch Pulverisette 6e planetary ball mill. Ball milling was carried out for different times up to three days in a WC–6 wt% Co system composed of a vial and of 16 balls of 10 mm diameter. A ball-to-powder weight ratio of 16:1 and a 400 rpm speed were adopted. The loading
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and unloading of the powders into and from the containers were always done inside an argon filled glove box. The powder blends were axially compacted at room temperature and under air with a pressure of 1.2 GPa. The pellets were subsequently sintered under vacuum (<10 2 Pa) by using a 300 °C/h heating rate followed by a 4 h heat treatment at 1800 °C completed by a 500 °C/h cooling stage. Laser diffractometry experiments were conducted on the powders to determine their size. Chemical analyses of the powders was performed, on the one hand by an inductively coupled plasma (ICP) atomic emission spectrometer to measure the cobalt content, and on the other hand by burn up using ELTRA ON-900e and LECO CS244e devices so as to evaluate carbon and oxygen contents. The microstructure, i.e. the particle morphology, the nature and the distribution of the second phases, the contamination by the milling system, was more especially studied by means of field emission gun/scanning electron microscopy (FEG/SEM), X-ray diffractometry (XRD) with a cobalt anticathode, electron probe microanalysis (EPMA) and transmission electron microscopy (TEM/STEM) analyses. The lattice strain and crystallites size were estimated from XRD results, applying the Warren–Averbach method [12] and using the {1 1 0} tungsten Bragg peaks. In the present paper, the word ÔcrystalliteÕ is going to designate a coherent diffraction domain within a single particle of powder whereas the term ÔgrainÕ is reserved for the sintered samples, where it keeps its usual meaning within bulk samples. The densities of the bulk sintered samples were measured by the ArchimedeÕs water immersion method while their Vickers microhardness was obtained under a 100 g load.
3. Results 3.1. Microstructure 3.1.1. Morphology and size of the powder particles Fig. 1a–c are scanning electron micrographs of the asreceived W and the W–1 vol%Y powder milled for 20 h,
Fig. 2. Laser diffractometry analysis of the particle size distribution of the W–1%Y powders milled for different times.
respectively. The third micrograph is typical of the case of a long-milled powder. The elemental tungsten powder is characterized by clusters of polyhedral particles with an average size of 4.5 lm (Fig. 1a). Such a morphology is consistent with tungsten powder arising from the reduction of WO3 by hydrogen [13]. Fig. 1c proves that ball milling has significantly changed the shape of the powder particles which becomes more equiaxed at long milling times. At the same time, the particle size is significantly reduced. Fig. 2 further shows that after 5 min of milling, welding seems to predominate whereas at the longest milling times (20 h), the mean particle size is finer and becomes submicrometric. It can be guessed, from such an observation, that fracture is the major mechanism at long milling times. This evolution of particle size is confirmed by bidimensional measurements of the powder particles on scanning electron micrographs [14,15]. Furthermore, it should be underlined that the milling time does not exert the same effect on the evolution of the size of the particles and of the crystallites (Section 3.1.2). 3.1.2. Inner microstructure of the powder particles and of the sintered samples Fig. 3a and b indicate the effect of milling time on the crystallite size and on the lattice strain of the samples measured along the [1 1 0] direction, respectively.
Fig. 1. Morphology of the elemental tungsten powder (a and b) and of the W–1%Y blend milled for 20 h, (c) BSE/FEG/SEM.
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Fig. 3. Effect of milling time on crystallite size (a) or on lattice strain (b) (XRD results).
For the powders, the longer the milling time, the finer the crystallite size is, up to a constant value. The present evolution suggests that the steady state of the process of ball milling is certainly attained from about 30 h of milling. Moreover it is worthy to note that, in this condition, the mean crystallite size is close to 5 nm. After sintering and at a given milling time, the full width at half maximum of the (1 1 0) peaks of tungsten diminishes, which indicates a growth of the grains during sintering (Fig. 3a). However the XRD results markedly differ from the measurements supplied from the FEG/SEM micrographs. For instance, Fig. 8c displays that the average grain size of an initially milled for 10 h sample grows to 10 lm after sintering, which is quite different from the value of 70 nm indicated by XRD results. Such a difference could be linked with two facts. Firstly, the XRD measurements are actually made in a preferential direction. Such a remark seems however to be of second importance because of the equiaxiality of the grains in the sintered samples. A second and more likely explanation lies in the fact that MurakamiÕs reagent etching only exhibits highly misoriented grain boundaries. Further works are necessary so as to be more conclusive on this topic. Contrary to the crystallite size, the lattice strain of the powder particles globally rises with the milling time to finally tend towards 0.3% (Fig. 3b). Ball milling generates many defects in the powder particles, which explains the growth of the average lattice strain. Nevertheless and in accordance with the literature data, the value of this lattice strain is always rather low and under its 0.58% estimated for pure milled tungsten [16]. The comparison with this upper result is not easy owing to different milling conditions.
phase on the X-ray diffractograms seems to arise more from both the too low nominal yttrium content and the marked absorption of the yttrium or yttria X-rays reflections by tungsten than from the fineness of the yttrium phase. Indeed, for the shortest milling times, the distribution of yttrium rich particles is rather coarse i.e. their spacing and their size are of about 7 and 3 lm, respectively. Moreover, according to the observations made with back scattered electrons (BSE)/FEG/SEM, these yttrium rich particles are only visible after the milling times shorter than 80 min. At least, in the latter conditions, they behave as less ductile particles than the tungsten ones which are flattened against their surface (Fig. 4). In the same way and as shown by Fig. 5, the EPMA results display a discrete distribution of yttrium rich phase within the particles of the W–1 vol%Y blend milled for 80 min. Moreover, it is significant to note that the same figure shows that such particles are not always oxidized. Such an observation also agrees with the lower than 624 ppm oxygen content of the powders milled for less than 40 min (Fig. 6c). The latter value is required in order to entirely change yttrium into yttria. The bulk metallic or Y(O) solid solution nature of this not oxidized yttrium can however be questioned. Otherwise, according to Fig. 6c, it is possible that yttrium is
3.1.3. Second phases This paragraph deals with the nature and the dispersion of the various kinds of second phases, including the yttrium rich ones. 3.1.3.1. Case of the blended powders. For any milling time, the absence of detection of an yttrium rich second
Fig. 4. W–17%Y powder milled for 80 min (BSE/FEG/SEM). Some white tungsten flakes can be seen over the surface of the grey yttrium rich particle.
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Fig. 5. Weight contents of oxygen, yttrium and tungsten within a large powder particle of W–1%Y milled for 80 min: (a) secondary electron image (the arrow indicates the position of the line of analysis) and (b) profiles of chemical composition.
Fig. 6. Effect of milling time on the contamination in mass percent of the powders by carbon (a), cobalt (b) and oxygen (c).
completely oxidized at the longest milling times. By way of contrast, the EPMA study failed to detect an yttrium rich phase in the W–1 vol%Y powders milled for 30 h. Such a result suggests that the yttrium rich phase is very fine in the final stationary stage of ball milling. The contamination of the blends has further been quantified. The chemical analyses show that the carbon and the cobalt contents of the blends continuously increase with the milling time (Fig. 6a and b). Moreover, and a priori in accordance with this observation because tungsten carbide and cobalt are the constituents of the milling system, Fig. 7a depicts a marked increase of the volume fraction of WC in the same powder for longer than 4 h milling times. These estimations were deduced from the measurements of the areas of the WC X-ray diffraction peaks (Fig. 7b) supposing both the complete transformation of yttrium into yttria and a value of the WC/Co ratio that is the same as the composition of the milling media. 3.1.3.2. Case of the sintered samples. Fig. 8a and b show that the longer the milling time, the more homogeneous and finer the second phases distribution. For any milling time, the size of the yttrium rich particles is however rather variable. For instance, and as depicted in Figs. 8c and 9, it is comprised between a few tens of nanometers and a few micrometers after the sintering of the blends milled for 20 h. Fig. 8c further shows that the second phases are essentially located at the grain boundaries. The EPMA analyses of the W–1 vol%Y sample sintered from the milled for 30 h powder
Fig. 7. (a) Contamination of the W–1%Y powders by the WC/Co milling system (XRD) and (b) example of X-ray diffraction pattern of the W–1%Y powder blend milled for 20 h.
mixture (Fig. 10) indicate that these second phases possess a stoichiometry close to yttria. The results are checked by the electron diffraction patterns of some nanometric white dispersoids like those shown in Fig. 9.
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Fig. 8. Dispersion of second phases in the W–1%Y sintered sample obtained from the blend milled for 40 min (a), 20 h (b) or 10 h (c). The latter sample was previously etched by the Murakami reagent in the latest case (BSE/FEG/SEM).
Fig. 9. STEM dark field of the yttria dispersoids in the W–1%Y sample sintered from the blend milled for 20 h.
Their crystal lattice is actually suited to such a chemical composition. Concerning the contamination of the sample sintered from the powder milled for 30 h, the STEM/EDX analyses indicate that the cobalt content is equal to 1.5 wt% in average in the matrix while it can attain values of 9.4 wt% in cobalt rich grains. 3.1.4. Properties of densified samples Fig. 11 shows that the green density is reduced from about 70% to 51% with the milling time. Moreover, up to 80 min of milling, the green density and the sintered density decrease in the same way while,
Fig. 10. EPMA analysis of the nature of an yttrium rich second phase in the W–1%Y sintered sample made of the milled for 30 h powder. The arrow indicates the position of the line of analysis.
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Fig. 11. Effect of milling time on both green and sintered density.
at longer milling times, the sintered density is improved contrary to the green density.
4. Discussion 4.1. Mechanism of ball milling In case of a ductile/ductile behavior of powders, the particles are at first flattened. This early stage is followed by a period dominated by particle welding which gives rise to lamellar particles. Afterwards the particles become more equiaxed, with first an oriented lamellar morphology, and subsequently a random arrangement of several lamellar granules within a given particle. Finally, a steady state is reached. It is characterized by a refinement of the inner microstructure while the size distribution of the particles remains constant [17,18]. The present evolution of the size and of the morphology of the particles agrees with Figs. 1 and 2 while it is more detailed in another paper [15] that clearly establishes that the W–1 vol%Y blend behaves as a ductile/ductile system during milling. 4.2. Powder microstructure The XRD 5 nm size of the crystallites inside the powder particles milled for the longest durations agrees with the literature data. At the steady state, a mean size of 10 nm has indeed been measured for pure tungsten [16]. Such an agreement is observed in spite of changes in the milling conditions and in the initial size and purity of the powder mixture. Moreover, such values correspond to the 4.5, 5 [16,19] and 9 nm [20] theoretical estimations of dislocation free crystallites in pure tungsten. Concerning the contamination, the carbon content accounts for both the carbon adsorption over the powder surfaces and the contamination of the powders by the milling system which is made of WC/Co. Figs. 6
and 7 prove that the latter contamination is very important. It very likely arises from an adhesive wear of the milling media because of both the tendency of cobalt and tungsten carbide to adhere [21] and the overall ductile behavior of the powder blends during milling [22]. Moreover, the values of cobalt and carbon obtained by chemical analyses agree with the results of X-ray diffractometry concerning the tungsten carbide content (Figs. 6 and 7). It proves that the carbon content in the powders is essentially supplied by the tungsten carbide of the milling system. The cobalt content measured by chemical analyses in the W–1 vol%Y sintered sample agrees with the decrease of the tungsten lattice parameter measured by XRD in the original blend to which decreases from 3.1648 A 3.1559 A after 30 h of milling. Such a diminution agrees with its estimate by using both the VegardÕs law and the chemical determination by burning up of the cobalt content in the powder. According to Fig. 6, the oxygen content increases with the milling time to attain very large values paying regard to the argon composition of the milling atmosphere. The origin of this oxygen and of its large content must therefore be questioned and the more as oxygen embrittles the tungsten [1]. The reactivity of the powders towards the traces of oxygen present within the milling atmosphere is very likely promoted by the fineness of the powder particles. However, the vial is not perfectly sealed as it has been checked by changing the chemical composition of the milling atmosphere that oxygen arises essentially from a leak. Likewise it is interesting to note that the oxygen content increases with the milling time. Many second phase particles have an effect on the plasticity of the blend. Firstly, the yttrium rich phases seem less ductile than the tungsten powder at least at early milling stages (Fig. 4), which is contrary to the literature data which mention a far higher ductility for the yttrium metal than for the tungsten one [23]. Such an anomaly is very likely to arise from the chemical composition of the yttrium rich particles, i.e. are they metallic or oxidized? According to both the present results and previous ones [15], a prolongation of the milling time promotes the fracture of tungsten particles, and increases the contamination by cobalt [which is metallic or oxidized (XPS experiments)] and tungsten carbide (Figs. 6b and 7a). Moreover, it can be suggested that for the longest durations of milling, yttrium is changed into yttria. The fracture of the tungsten particles could be explained by an enhancement of their plastic deformation and therefore of their work hardening that entails their fracture. Plastic deformation would then be favored by the stress concentration due to the hard particles of yttria and of tungsten carbide. According to its high volume fraction (Fig. 7), tungsten carbide plays a dominant
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role on the general behavior of the powder. In addition, the formation of a cobalt solid solution in the tungsten is likely to modify the mechanical behavior of the blend. 4.3. Densified samples The decrease of the green density with the milling time can be due to many factors. The green density actually depends on a great number of parameters dealing with the morphology, the size distribution and the ductility of the powder particles. Such a ductility is governed by both the distribution of the second phases within the compacts and the degree of contamination and the dislocation substructure of the powders. According to the present results concerning the density of the green compacts and the properties of the W–1 vol%Y milled powders, two assumptions may be considered. Firstly, the refinement of the particles size with the milling time gives rise to a narrow monomodal size distribution (Fig. 2) which is less favorable than a bimodal size distribution to eliminate the porosity within both a green compact [16] and a sintered sample [24]. Secondly, the rise of the contamination, for instance by oxygen (Fig. 6c), can impede the particles to weld during the stage of compaction. The oxide layer on the powder particles surfaces prevents them to deform and therefore contributes to a low densification. It is however not possible to check the latter hypothesis because of the easiness of the green compacts to crumble. Over 4 h of milling, the sintered density rises whereas the green density continues to decrease (Fig. 11). The origin of such a discrepancy is very likely to have to do with the nature of the sintering mechanism that can however be questioned. Some clues suggest that the sintering proceeds in solid phase whereas others could indicate that it is promoted by the formation of a liquid phase. In the following, both cases are successively considered. Solid phase sintering seems to occur because of the presence of polyhedral grains (Fig. 8b and c). Solid phase sintering is especially plausible because the sintering temperature is rather low for tungsten, which is usually solid phase sintered at temperatures around 2000 °C [24]. Moreover, the contamination by the milling system, i.e. by cobalt and tungsten carbide, is obvious from 4 h of milling, which suggests that cobalt has to do with the rise of the sintered density that occurs in spite of the decrease of the compactÕs one. The literature data actually indicate that the sintering of tungsten can be activated by the presence of cobalt in solid solution because this chemical element entails an increase of the diffusion rate [25–27]. Moreover, cobalt is actually in solid solution in the sintered sample insofar as the 1.5 wt% content of cobalt found by STEM/EDS in the W–1 vol%Y milled for 30 h and sintered sample agrees with the VegardÕs law.
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Alternatively, various features suggest that the sintering of tungsten can locally proceed in liquid phase. Firstly, the liquid phase may arise from the melting of isolated islands of the WC/Co cermet which is usually liquid phase sintered at about 1450 °C [27,28]. Secondly, pure yttrium and a 0–0.304 at% O–Y solid solution melt at 1522 °C [29] and 1522–1670 °C [30], respectively, that is to say at temperatures that are significantly lower than the 1800 °C present sintering temperature. The latter consideration agrees with both the presence of some rounded interfaces with the second phases (Fig. 8a) as well as with the growth of intergranular second phases during sintering. Such a growth has been noticed by the present authors comparing the microstructures of sintered samples of W plus 1 vol% or 17 vol%Y [15]. Itoh et al. [31] have also noticed the growth of the yttria particles and its increase with both the yttria content and the sintering temperature in W– Y2 O3 samples. Due to the very low solubility of yttrium in pure tungsten [29], such phenomena are rather astonishing and the more astonishing as, in the present case, the presence of cobalt cannot improve the solubility because of the tendency of yttrium and cobalt to form intermetallic compounds [32]. Thirdly, liquid phase sintering can further arise from the formation either of an 0.0785 at% W–Y eutectic at 1522 °C [31] or a W–Y2 O3 one at 1560 °C [30]. At last, liquid phase sintering may again proceed because of the formation of a Y2 (WO4 )3 phase which melts at 1440 °C. Such a compound has indeed been mentioned by Itoh et al. in W–Y2 O3 blends [31]. Consequently, numerous sintering mechanisms are very likely to intervene in a more or less exclusive way. Some further experiments would be necessary in order to be more conclusive. Otherwise, the formation of finer and finer nanostructures with the milling time usually improves the densification [9]. The comparison between Figs. 3 and 11 however indicates that this parameter is not predominant owing to a successive decline and rise of the sintered density with the prolongation of milling. At last, Fig. 3a, one more time, confirms that a significant grain growth is usual during the sintering of nanostructures (Section 3.1.2). Such a growth of the crystallites during sintering is however the less important, the longer the milling time (Fig. 3a) i.e. as the sizes of the particles and of the crystallites inside these particles decrease with the milling time. The whole of these observations suggest that the content and the distribution of the second phases are not suitable in order to fully stabilize a nanometric crystallite size of the powders by the Zener pinning mechanisms of grain boundaries [31,33] (Sections 3.1.3 and 3.1.4). However, after the longest milling times, the quality of the second phases distribution is more convenient than after the first steps of milling.
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4.4. Prospects The use of a WC/Co milling system brings about a high contamination of the samples by tungsten carbide and cobalt for the longest milling times, which could be avoided by changing the nature of the milling system. Such a contamination nevertheless presents an advantage as cobalt may improve the density of the sintered alloys (Section 4.3). Likewise an eventual change of the nature of the milling system could be questioned in order to maintain the stability of the nanostructure at the longest milling time because the sole yttrium content could then be insufficient to limit the grain growth. Moreover, it is worthy to control the milling atmosphere so as to avoid or at least to limit the contamination by oxygen and then to satisfy the proposed goal.
5. Summary and conclusions 1. The present study prominently establishes the interest of long milling times in order to elaborate ODS alloys and nanostructures, which is confirmed by many authors [34]. A prolongation of the milling duration generates firstly a refinement of the size of both the powder particles and of their crystallites and secondly an improved distribution of the yttrium rich phase (and of the second phases in general). At the longest milling time, it can be guessed that the second phases are suitably dispersed so as to impede the displacement of grain boundaries: the increase of the crystallite size measured by XRD actually remains small after sintering. 2. The use of a WC/Co milling system brings about a high contamination of the samples by tungsten carbide and cobalt for the longest milling times. 3. Otherwise, the high chemical affinity of the powders towards oxygen remains to be solved. 4. The microstructural analyses of the sintered samples can validate a solid phase sintering mechanism with local liquid phase sintering. The occurrence of liquid phase sintering is probably the cause of the opposite evolutions of the green compact and the sintered densities.
Acknowledgements The authors are very grateful to N. Llorca (CEA Saclay, France) for giving them the opportunity of using a gloves box, to C. Boyaval (IEMN, Lille, France) for the FEG/SEM micrographs, to J.F. Lartigue (CERMEP Grenoble, France) for the chemical analyses of the powders, to O. Duplessis (CEA Saclay, France) for the laser diffractometry analyses, to N. Lochet (CEA Saclay, France) for the sintering, to C. Fucili (CEA Valrh^ o,
France) for the EPMA analyses and to L. Gengembre (University of Lille, Laboratory of Catalysis, France) for the XPS analyses.
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