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Thermal stability of metastable Mg–30% Ti–2% Al–0.9% Mn (wt.%) alloy synthesised by PVD
Journal of Alloys and Compounds 387 (2005) 115–120
Thermal stability of metastable Mg–30% Ti–2% Al–0.9% Mn (wt.%) alloy synthesised by PVD G. Garc´es∗ , P. P´erez, P. Adeva Department of Physical Metallurgy, National Centre for Metallurgical Research (CENIM–CSIC), Av. De Gregorio del Amo 8, 28040 Madrid, Spain Received 29 April 2004; received in revised form 1 June 2004; accepted 1 June 2004
Abstract The objective of this work is the study of the thermal stability of the Mg–30% Ti–2% Al–0.9% Mn (wt.%) alloy produced by physical vapour deposition (PVD). For this purpose differential scanning calorimetry and transmission electron microscopy techniques have been used. The alloy, in the as-deposited condition, is a solid solution of the alloying elements in the magnesium matrix, which induces a strong decrease in lattice parameters. The microstructure is characterised by elongated grains oriented in the deposit growth direction. After a differential scanning calorimetry experiment, the breakdown of the solid solution occurs. The precipitation phenomenon takes place by three exothermal reactions. As in Mg–Ti alloys, the two first peaks are associated with the titanium precipitation within the grains. The third transformation is due to the precipitation at grain boundaries of spherical Ti3 Al particles. At higher temperatures, above 720 K, an increase in heat flow takes place and it is related to the strong oxidation of the magnesium matrix. © 2004 Elsevier B.V. All rights reserved. Keywords: PVD growth; Magnesium alloy; Precipitation; Kinetics
1. Introduction Magnesium alloys present a great potential as structural materials in the aerospace and automobile industries mainly because of their low densities and high specific resistance. Their main problem is the poor corrosion resistance. At the present, it is a dedicating great effort to investigate different alternatives to improve the corrosion behaviour. One of them consists of the addition of alloying elements that oxidise in preference to magnesium or can build an oxide scale when magnesium is depleted. The low solubility in molten magnesium of interesting alloying elements, such as titanium, vanadium, chromium, etc. makes it necessary to study new techniques to synthesise new corrosion-resistant magnesium alloys. Physical vapour deposition (PVD) is a non-equilibrium solidification method very promising to develop these kinds of alloys. Although this technique is usually associated with the ∗
Corresponding author. Tel.: +34-1-5538900; fax: +34-1-5347425. E-mail address: [email protected] (G. Garc´es).
0925-8388/$ – see front matter © 2004 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2004.06.029
production of thin films or coatings, DRA Farnborough developed an experimental system to produce different aluminium and magnesium alloys of several millimetres of thickness [1–6]. The results have shown that this process led to a refinement of the microstructure and an extension of the solid solubility limits of the alloying elements. The main advantage of the PVD is that the alloying elements do not need to be melted together to produce the alloy and therefore problems of solubility in molten magnesium are overcome. Some studies about the improvement of corrosion resistance of magnesium alloys indicate that titanium additions higher than 20 wt.% reduce the corrosion rate of PVD magnesium in chloride solutions. It is proposed that titanium stabilises the surface film on the alloy surface. Baliga et al. have associated this improvement with a rutile film formation, which provide protection from attack by chlorine ions in the saline environment where dissolution of magnesium occurs [7]. The solid solution breakdown temperature is an important parameter to evaluate the temperature range of using in this kind of metastable alloys. Studies about microstructure and thermal stability of PVD Mg–Ti alloys indicate that the
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solid solution breakdown occurs by the precipitation of pure titanium taking place in two stages. At the first stage, titanium plates precipitate in the basal planes, which grow for higher temperatures in the [0 0 0 1] direction, by a process of renucleation above the pre-existent plates [8–10]. The present work is focused on the study of the solid solution breakdown temperature of a PVD Mg–30% Ti–2% Al–0.9% Mn (wt.%) alloy during a continuous heating. The precipitation sequence is followed by transmission electron microscopy (TEM) and the effect of the presence of aluminium and manganese on the precipitation reaction is compared with that of binary Mg–Ti alloys.
2. Experimental procedure The magnesium alloy, Mg–14% Ti–1% Al–0.9% Mn (wt.%) of nominal composition, was grown by Qinetiq using the PVD technique with an aluminium collector at ≈182 ◦ C. Details about the synthesis procedure can be consulted in reference [1]. A marked chemical inhomogeneity consisting of a sharp decrease in the titanium content in the middle of the deposit was observed. This study is carried out in the initial zone next to the collector surface with a thickness of 750 m where an average composition of Mg–30% Ti–2% Al–0.9% Mn (wt.%) was measured. In this zone a stable presence of titanium, aluminium and manganese is insured. Microstructural characterisation of the alloy in the asdeposited condition was carried out through X-ray diffraction as well as optical and electron microscopy. The precipitation phenomenon was followed utilising DSC and TEM. X-ray diffraction was used to characterise the alloy in the as-deposited condition as well as after DSC scan and for measuring lattice parameters. The analysis was carried out on a D5000 SIEMENSTM diffractometer using -filtered Cu K␣ radiation. Metallographical preparation of the scanning samples consisted in mechanical polishing, and etching in a solution of 10% Keller, 40% ethylene glycol and 50% water. Specimens for transmission electron microscopy were prepared by electrolytic polishing using a mixture of 3% percloric acid, 33% butanol, 64% methanol. Ion milling at liquid nitrogen temperature was used to remove a fine oxide film formed in the surface. Then, the samples were examined in a JEOL JEM-2032 microscope operating at 200 kV. DSC measurements were carried out in a Perkin-Elmer System 4 Thermal Analyser. The scans were made under argon atmosphere to minimise oxidation at heating rates of 10, 20, 30 and 40 K min−1 . High purity cast magnesium was used as reference material. To obtain a baseline (which depends on, among other factors, the different heat capacities of the reference and sample pans) two runs were carried out. The first run was conducted using high purity magnesium discs of approximately the same mass in both pans. The second run was carried out after replacing the magnesium disc in the sample pan by a disc of the material under study. Subtrac-
tion of the heat flow in the first run from the heat flow in the second run results in a signal, which allows the calculation of the heat flow due to reactions in the sample. The different states of precipitation were studied stopping the DSC scan at the peak temperatures and quenching the samples.
3. Results and discussion The X-ray diffraction pattern of the alloy in the asdeposited condition taken in the plane perpendicular to the deposit growth direction is shown in Fig. 1a deviation from the theoretical peak for pure magnesium can be observed. The alloy lattice parameters calculated from the X-ray diffraction pattern are 0.31596 and 0.50812 nm, respectively. A strong decrease in lattice parameters, with respect to those of pure magnesium is noticed in Fig. 2 which compares the lattice parameter values of the alloy with respect to those obtained from bibliography in the Mg–Ti binary system. Thus, it can be followed that the alloying elements, especially the titanium atoms, provoke a strong contraction in the magnesium lattice. The microstructure throughout the thickness consists of columnar grains oriented in the deposit growth direction
Fig. 1. X-ray diffraction pattern of the alloy in the as-deposited condition. Lines show the position and intensity of the theoretical pure magnesium peaks.
Fig. 2. Lattice parameter of the Mg–Ti system as function of titanium concentration.
G. Garc´es et al. / Journal of Alloys and Compounds 387 (2005) 115–120
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Table 1 Values of the first exothermal peak temperature for each heating rate
Fig. 3. Bright field image of the columnar grain structure of the alloy in the as-deposited condition.
(Fig. 3). Since no precipitates are observed under the asdeposit condition and the lattice parameters are strongly decreased, it is assumed that the alloying elements are in solid solutions. The microstructure of columnar grains observed in this deposit indicates a growth process in the zone II of the Movchan and Demchishin structure zone model (SZM) developed for PVD materials [11,12]. According to this model the microstructure features of PVD alloys depend on the ratio Tc /Tm , where Tc and Tm are the collector and melting temperature, respectively. This ratio governs the diffusion mechanisms of adatoms in the surface of the alloy. Zone II is defined in the range of 0.3 < Tc /Tm < 0.5 and the surface diffusion of the adatom controls the deposit growth. The grain growth process (in the deposit growth direction) has been described in terms of surface recrystallization, evolutionary selection and granular epitaxy [11,13]. Although the alloy was grown in zone II, the substrate temperature is low enough (approximately 0.3) to produce the solid solution of the alloying elements in the magnesium matrix, provoking strong contraction of the lattice cell. The thermal stability of the alloy was followed by differential scanning calorimetry using different heating rates (Fig. 4). Three exothermal peaks can be observed, although only the first peak is clearly distinguished for all heating rate.
Fig. 4. Heat flow vs. temperature from the DSC experiment at 10, 20 and 30 K min−1 .
Heating rate (K min−1 )
Tfirst peak (K)
10 20 30 40
558 569 575 580
Table 1 shows the peak temperature for each heating rate in the case of the first transformations. As shown Fig. 4, these transformations occur at higher temperatures with increasing the heating rate as it is common in thermally activated transformations. The three transformations are associated with the solid solution breakdown. In previous works the thermal stability of PVD alloys deposited at lower temperatures (zone I in the SZM) was studied. A stress relaxation phenomenon due to the high dislocations density produced during the deposit growth was reported [8,14,15]. However, in the case of the alloy of this study, that reaction has not been present. This can be related to the higher deposition temperature used to grow this deposit, which induces surface diffusion of ad-atoms and therefore a lower density of defects will be generated. In order to follow the precipitation sequence, the microstructure of the alloy quenched at the first peak as well as at the end of the DSC scan was examined by TEM. Fig. 5a shows the TEM image of the sample quenched at the first peak. Small precipitates of around 2 nm perpendicular to the [0 0 0 1] direction are clearly observed when the g = (0 0 0 2) vector was excited. These precipitates are uniformly distributed in the magnesium matrix. Since the SAD pattern from the [1 1 2¯ 0] zone axis shows neither extra reflections nor diffuse scattering, it can be concluded that precipitates are fully coherent with the magnesium matrix. The precipitates are parallel to each other and perpendicular to the [0 0 0 1] direction. An elastic distortion, expected around precipitates with a value of around the 8%, is mostly adopted by a magnesium matrix [10]. Thus, an increase in the alloy hardness as the precipitates grow in the basal plane is expected. Fig. 6 shows the microhardness evolution during an isothermal treatment at 553 K. As can be seen, the alloy exhibits high hardness values under the as-deposited condition due to the solid solution reinforcement effect of alloying elements. As the thermal treatment progresses, the hardness increases due to the nucleation and growth of the titanium plates in the basal plane reaching a maximum after 30 h of annealing. Then, a slight decrease takes place due to coarsening of the titanium precipitates. The kinetic parameters of the process have been calculated to compare the precipitation reaction in this first peak with respect to the binary Mg–Ti alloy. The activation energies for the transformation can be determined by means of Kissinger’s analysis [16] as follows: Tp 2 E E ln (1) + ln ≈ RTp RK0 Φ
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Fig. 7. Kissinger plot of the first DSC precipitation peak.
tivation energies can be calculated. The values of E and K0 obtained from the Kissinger fit for the first are 160 kJ mol−1 and 10−13 s−1 , respectively, similar to those obtained in the binary system Mg–12%Ti [9]. Thus, there is no influence of the other alloying elements on the precipitation of pure titanium. On the other hand, the sample quenched at the end of the DSC scan shows coarser precipitates and extra double diffraction spots come out in the SAD pattern of the [1 1 2¯ 0] zone axis (Fig. 5b). The thickness of these precipitates in the [0 0 0 1] direction increases with respect to the initial precipitates (5–10 nm). Moreover, the precipitates are responsible for the appearance of Moir´e fringes when the (0 0 0 2) g vector is excited. In the case of the parallel Moir´e, the fringe spacing D of the structure precipitate-matrix is given by [17]: Fig. 5. (a–c) Bright field images and SADP from the [] zone axis showing the morphology of the titanium precipitates of the sample (a) quenched at first precipitation peak and (b) at the end of the DSC. Bright field image (c) and SADP (d) from the [0 0 1] zone axis of the microstructure of the sample quenched at the end of the DSC.
where E is the apparent activation energy, Tp the temperature of the maximum reaction rate, K0 the pre-exponential factor, R the universal gas constant and Φ the heating rate. The apparent activation energies for the transformation is obtained from the slope of the plot of ln (Tp 2 /Φ) versus 1/ Tp as shown in Fig. 7 for the case of the first peak where the temperature of the maximum is clearly distinguished. This figure shows a good linear correlation and the values of the apparent ac-
Fig. 6. Microhardness evolution during an isothermal treatment at 553 K.
1 1 1 = − D d1 d2
(2)
where d1 and d2 are the interplanar distance of the two superimposed crystal lattices. From Fig. 5b, the values for D can be measured obtaining an average value of 2.65 nm. Therefore, it is possible to calculate the interplanar distance of the precipitates taking as the interplanar distance between the (0 0 0 2) magnesium planes its theoretical value (0.260 nm). The value obtained for d2 is 0.236 and it is in concord with the theoretical interplanar distance between the (0 0 0 2) ␣-titanium planes. The epitaxial relationship of these titanium precipitates is expressed by [0 0 0 1]Ti ||[0 0 0 1] Mg and [1 1 2¯ 0]Ti ||[1 1 2¯ 0] Mg according with previous works [8–10]. It is worthy to note that the titanium precipitates present a hexagonal feature with their face parallel to the {1 1 2¯ 0} planes when the sample is observed in the [0 0 0 1] zone axis (Fig. 5c) and ellipsoidal with faces parallel to the traces of the {0 0 0 1} and {0 1 1¯ 1} planes when the sample is observed in the [1 1 2¯ 0] zone axis (Fig. 5b). During the precipitate coarsening the coherence between the titanium particles and the magnesium matrix is lost, provoking the appearance of misfit dislocation between the precipitates as it is observed in the weak beam image in Fig. 8.
G. Garc´es et al. / Journal of Alloys and Compounds 387 (2005) 115–120
Fig. 8. Weak beam image of the microstructure of the sample at the end of the DSC scan.
During the non-isothermal treatment, precipitation of the alloying elements takes place. The first two exothermal peaks are due to titanium precipitation, which takes place in two sequences as reported in binary Mg–Ti alloys. In the first peak, it can be supposed that titanium precipitates initially nucleate and grow to occupy a certain area in the (0 0 0 1) planes. The presence of small precipitates, which were observed when the (0 0 0 2) g vector was excited supports this argument. When the increase in misfit energy associated with small precipitate
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(basal plane) growth cannot be accomodated by the matrix, these precipitates grow along the [0 0 0 1] direction loosing their coherence with the matrix (second peak). Hitherto, it is assumed that this precipitates was incoherent or semicoherent with the matrix. However, the appearance of misfit dislocations generated in the precipitate interfaces has not been reported. The appearance of these misfit dislocations (Fig. 8) supports the assessment that these precipitates are incoherent. Not only the titanium precipitates are present at this stage but also spherical particles of about 20 nm decorating the grain boundaries are visible as shown in Fig. 9. The EDS results indicate that these particles are titanium and aluminium rich. The precipitation at grain boundaries of spherical Ti3 Al precipitates above 750 K, is consistent with the ternary equilibrium phase diagram of Mg–Al–Ti [18]. The isothermal section above 700 K shows that an alloy of similar composition as the alloy of this work would be composed of ␣Ti and the ␣2 phase (Ti3 Al DO19 ) in a magnesium matrix. This observation also agrees with recently studies of Wang et al. [19] where magnesium reinforced by in situ TiC particles was synthesized. Titanium, carbon and aluminium powders and Mg–9% Al–1% Zn (wt.%) alloy were used as the reinforced phase precursor and the matrix, respectively. During the composite processing not only TiC particles but also Ti3 Al precipitates was formed.
Fig. 9. Bright field image showing spherical precipitates at the grain boundaries. EDS spectra of the Mg matrix, spherical precipitates at grain boundaries and Ti precipitates within the grain.
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Finally, the continuous increase in the heat flow above 703 K of Fig. 3 is connected with the oxidation of the deposit [3,8]. The high concentration of defects, especially pores at the grain boundaries, induce rapid inward oxygen transport provoking the increase in the heat flow and mass gain in these PVD thick deposits [20,21].
4. Summary The thermal stability and the precipitation reaction in the PVD Mg–30% Ti–2% Al–0.9% Mn (wt.%) alloy have been studied. The microstructure of the alloy throughout the thickness consists of columnar grains in the deposit growth direction constituted by a solid solution of alloying elements in the magnesium matrix that provokes a strong contraction of the magnesium lattice parameters. During DSC experiments three exothermal transformation that correspond to precipitation phenomena in the magnesium matrix have been observed. The precipitation of titanium occurs like in the binary Mg–Ti alloy, in two steps. Furthermore, the presence of aluminium leads to the formation of the intermetallic phase Ti3 Al at the grain boundaries.
Acknowledgements The authors would like to thank Dr. S.B. Dodd, S. Morris and R.C. Piller of DERA (Qinetiq, actually), Farnborough, for providing the vapour deposited alloys along with useful information on processing conditions. We gratefully acknowledge the support of the Comisi´on Interministerial de Ciencia y Tecnolog´ıa (MAT 981620-CE) and Comunidad de Madrid (CAM 700-739).
References [1] S.B. Dodd, R.W. Gardiner, in: Proceedings of the Third International Magnesium Conference, The Institute of Metals, Manchester, 1996, p. 271. [2] S.B. Dodd, S. Morris, R.W. Gardiner, in: Proceedings of the Magnesium Alloys and their Applications, Frankfurt: WerkstoffInformationgesellschaft, 1998, p. 375. [3] S. Diplas, P. Tsakiropoulos, R.M.D. Brydson, Mater. Sci. Technol. 14 (1998) 689–698. [4] K.R. Baldwin, D.J. Bray, G.D. Howard, R.W. Gardiner, Mater. Sci. Technol. 12 (1996) 937–943. [5] S. Morris, S.B. Dodd, P.J. Hall, A.J. Mackinnom, L.E. Berlouis, J. Alloys Compd. 293–295 (1999) 458–462. [6] P. Adeva, S.B. Dodd, P. Morgan, F. Hehmann, P. Steinmetz, F. Sommer, Adv. Eng. Mater. 3 (2001) 147–152. [7] C.B. Baliga, P. Tsakiropoulos, S.B. Dodd, R.W. Gardiner, in: Proceedings of the Third International Magnesium Conference, The Institute of Metals, Manchester, 1996, p. 627. [8] K.E. Bagnall, P.G. Partridge, J.W. Steed, S.B. Dodd, R.W. Gardiner, in: Proceedings of the Third International Magnesium Conference, The Institute of Metals, Manchester, 1996, p. 299. [9] G. Garc´es, P. Adeva, Phil. Mag. A 82 (2002) 699–714. [10] G. Garc´es, P. P´erez, P. Adeva, Scripta Mater. 45 (2001) 1001–1007. [11] B.A. Movchan, A.V. Demchishin, Phys. Metal. Metallogr. 28 (1969) 653–660. [12] J.A. Thornton, Rev. Mater. Sci. 7 (1977) 239–260. [13] C.R.M. Grovenor, H.T.G. Hertzell, D.A. Smith, Acta Metall. 32 (1984) 773–781. [14] S. Diplas, P. Tsakiropoulos, R.M.D. Brydson, Mater. Sci. Technol. 15 (1999) 1359–1372. [15] J.P. Chu, C.H. Chung, P.Y. Lee, J.M. Rigsbee, J.Y. Wang, Metall. Mater. Trans. A 29A (1998) 647–658. [16] E.J. Mittemeijer, J. Mater. Sci. 27 (1992) 3977–3987. [17] D.B. Williams, C.B. Carter, Transmission Electron Microscopy, Plenum Press, New York, 1996. [18] G. Petzow, G. Effenberg, Ternary Alloys, vol. 6, ASM International, New York, Basel, 1993, p. 443. [19] H.Y. Wang, Q.C. Jiang, X.L. Li, J.G. Wang, Scripta Mater. 48 (2003) 1349–1354. [20] P. P´erez, G. Garc´es, P. Adeva, Oxid. Met. 58 (2002) 60. [21] P. P´erez, G. Garc´es, P. Adeva, Mater. Corr. 54 (2003) 145.
Thermal stability of metastable Mg–30% Ti–2% Al–0.9% Mn (wt.%) alloy synthesised by PVD G. Garc´es∗ , P. P´erez, P. Adeva Department of Physical Metallurgy, National Centre for Metallurgical Research (CENIM–CSIC), Av. De Gregorio del Amo 8, 28040 Madrid, Spain Received 29 April 2004; received in revised form 1 June 2004; accepted 1 June 2004
Abstract The objective of this work is the study of the thermal stability of the Mg–30% Ti–2% Al–0.9% Mn (wt.%) alloy produced by physical vapour deposition (PVD). For this purpose differential scanning calorimetry and transmission electron microscopy techniques have been used. The alloy, in the as-deposited condition, is a solid solution of the alloying elements in the magnesium matrix, which induces a strong decrease in lattice parameters. The microstructure is characterised by elongated grains oriented in the deposit growth direction. After a differential scanning calorimetry experiment, the breakdown of the solid solution occurs. The precipitation phenomenon takes place by three exothermal reactions. As in Mg–Ti alloys, the two first peaks are associated with the titanium precipitation within the grains. The third transformation is due to the precipitation at grain boundaries of spherical Ti3 Al particles. At higher temperatures, above 720 K, an increase in heat flow takes place and it is related to the strong oxidation of the magnesium matrix. © 2004 Elsevier B.V. All rights reserved. Keywords: PVD growth; Magnesium alloy; Precipitation; Kinetics
1. Introduction Magnesium alloys present a great potential as structural materials in the aerospace and automobile industries mainly because of their low densities and high specific resistance. Their main problem is the poor corrosion resistance. At the present, it is a dedicating great effort to investigate different alternatives to improve the corrosion behaviour. One of them consists of the addition of alloying elements that oxidise in preference to magnesium or can build an oxide scale when magnesium is depleted. The low solubility in molten magnesium of interesting alloying elements, such as titanium, vanadium, chromium, etc. makes it necessary to study new techniques to synthesise new corrosion-resistant magnesium alloys. Physical vapour deposition (PVD) is a non-equilibrium solidification method very promising to develop these kinds of alloys. Although this technique is usually associated with the ∗
Corresponding author. Tel.: +34-1-5538900; fax: +34-1-5347425. E-mail address: [email protected] (G. Garc´es).
0925-8388/$ – see front matter © 2004 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2004.06.029
production of thin films or coatings, DRA Farnborough developed an experimental system to produce different aluminium and magnesium alloys of several millimetres of thickness [1–6]. The results have shown that this process led to a refinement of the microstructure and an extension of the solid solubility limits of the alloying elements. The main advantage of the PVD is that the alloying elements do not need to be melted together to produce the alloy and therefore problems of solubility in molten magnesium are overcome. Some studies about the improvement of corrosion resistance of magnesium alloys indicate that titanium additions higher than 20 wt.% reduce the corrosion rate of PVD magnesium in chloride solutions. It is proposed that titanium stabilises the surface film on the alloy surface. Baliga et al. have associated this improvement with a rutile film formation, which provide protection from attack by chlorine ions in the saline environment where dissolution of magnesium occurs [7]. The solid solution breakdown temperature is an important parameter to evaluate the temperature range of using in this kind of metastable alloys. Studies about microstructure and thermal stability of PVD Mg–Ti alloys indicate that the
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solid solution breakdown occurs by the precipitation of pure titanium taking place in two stages. At the first stage, titanium plates precipitate in the basal planes, which grow for higher temperatures in the [0 0 0 1] direction, by a process of renucleation above the pre-existent plates [8–10]. The present work is focused on the study of the solid solution breakdown temperature of a PVD Mg–30% Ti–2% Al–0.9% Mn (wt.%) alloy during a continuous heating. The precipitation sequence is followed by transmission electron microscopy (TEM) and the effect of the presence of aluminium and manganese on the precipitation reaction is compared with that of binary Mg–Ti alloys.
2. Experimental procedure The magnesium alloy, Mg–14% Ti–1% Al–0.9% Mn (wt.%) of nominal composition, was grown by Qinetiq using the PVD technique with an aluminium collector at ≈182 ◦ C. Details about the synthesis procedure can be consulted in reference [1]. A marked chemical inhomogeneity consisting of a sharp decrease in the titanium content in the middle of the deposit was observed. This study is carried out in the initial zone next to the collector surface with a thickness of 750 m where an average composition of Mg–30% Ti–2% Al–0.9% Mn (wt.%) was measured. In this zone a stable presence of titanium, aluminium and manganese is insured. Microstructural characterisation of the alloy in the asdeposited condition was carried out through X-ray diffraction as well as optical and electron microscopy. The precipitation phenomenon was followed utilising DSC and TEM. X-ray diffraction was used to characterise the alloy in the as-deposited condition as well as after DSC scan and for measuring lattice parameters. The analysis was carried out on a D5000 SIEMENSTM diffractometer using -filtered Cu K␣ radiation. Metallographical preparation of the scanning samples consisted in mechanical polishing, and etching in a solution of 10% Keller, 40% ethylene glycol and 50% water. Specimens for transmission electron microscopy were prepared by electrolytic polishing using a mixture of 3% percloric acid, 33% butanol, 64% methanol. Ion milling at liquid nitrogen temperature was used to remove a fine oxide film formed in the surface. Then, the samples were examined in a JEOL JEM-2032 microscope operating at 200 kV. DSC measurements were carried out in a Perkin-Elmer System 4 Thermal Analyser. The scans were made under argon atmosphere to minimise oxidation at heating rates of 10, 20, 30 and 40 K min−1 . High purity cast magnesium was used as reference material. To obtain a baseline (which depends on, among other factors, the different heat capacities of the reference and sample pans) two runs were carried out. The first run was conducted using high purity magnesium discs of approximately the same mass in both pans. The second run was carried out after replacing the magnesium disc in the sample pan by a disc of the material under study. Subtrac-
tion of the heat flow in the first run from the heat flow in the second run results in a signal, which allows the calculation of the heat flow due to reactions in the sample. The different states of precipitation were studied stopping the DSC scan at the peak temperatures and quenching the samples.
3. Results and discussion The X-ray diffraction pattern of the alloy in the asdeposited condition taken in the plane perpendicular to the deposit growth direction is shown in Fig. 1a deviation from the theoretical peak for pure magnesium can be observed. The alloy lattice parameters calculated from the X-ray diffraction pattern are 0.31596 and 0.50812 nm, respectively. A strong decrease in lattice parameters, with respect to those of pure magnesium is noticed in Fig. 2 which compares the lattice parameter values of the alloy with respect to those obtained from bibliography in the Mg–Ti binary system. Thus, it can be followed that the alloying elements, especially the titanium atoms, provoke a strong contraction in the magnesium lattice. The microstructure throughout the thickness consists of columnar grains oriented in the deposit growth direction
Fig. 1. X-ray diffraction pattern of the alloy in the as-deposited condition. Lines show the position and intensity of the theoretical pure magnesium peaks.
Fig. 2. Lattice parameter of the Mg–Ti system as function of titanium concentration.
G. Garc´es et al. / Journal of Alloys and Compounds 387 (2005) 115–120
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Table 1 Values of the first exothermal peak temperature for each heating rate
Fig. 3. Bright field image of the columnar grain structure of the alloy in the as-deposited condition.
(Fig. 3). Since no precipitates are observed under the asdeposit condition and the lattice parameters are strongly decreased, it is assumed that the alloying elements are in solid solutions. The microstructure of columnar grains observed in this deposit indicates a growth process in the zone II of the Movchan and Demchishin structure zone model (SZM) developed for PVD materials [11,12]. According to this model the microstructure features of PVD alloys depend on the ratio Tc /Tm , where Tc and Tm are the collector and melting temperature, respectively. This ratio governs the diffusion mechanisms of adatoms in the surface of the alloy. Zone II is defined in the range of 0.3 < Tc /Tm < 0.5 and the surface diffusion of the adatom controls the deposit growth. The grain growth process (in the deposit growth direction) has been described in terms of surface recrystallization, evolutionary selection and granular epitaxy [11,13]. Although the alloy was grown in zone II, the substrate temperature is low enough (approximately 0.3) to produce the solid solution of the alloying elements in the magnesium matrix, provoking strong contraction of the lattice cell. The thermal stability of the alloy was followed by differential scanning calorimetry using different heating rates (Fig. 4). Three exothermal peaks can be observed, although only the first peak is clearly distinguished for all heating rate.
Fig. 4. Heat flow vs. temperature from the DSC experiment at 10, 20 and 30 K min−1 .
Heating rate (K min−1 )
Tfirst peak (K)
10 20 30 40
558 569 575 580
Table 1 shows the peak temperature for each heating rate in the case of the first transformations. As shown Fig. 4, these transformations occur at higher temperatures with increasing the heating rate as it is common in thermally activated transformations. The three transformations are associated with the solid solution breakdown. In previous works the thermal stability of PVD alloys deposited at lower temperatures (zone I in the SZM) was studied. A stress relaxation phenomenon due to the high dislocations density produced during the deposit growth was reported [8,14,15]. However, in the case of the alloy of this study, that reaction has not been present. This can be related to the higher deposition temperature used to grow this deposit, which induces surface diffusion of ad-atoms and therefore a lower density of defects will be generated. In order to follow the precipitation sequence, the microstructure of the alloy quenched at the first peak as well as at the end of the DSC scan was examined by TEM. Fig. 5a shows the TEM image of the sample quenched at the first peak. Small precipitates of around 2 nm perpendicular to the [0 0 0 1] direction are clearly observed when the g = (0 0 0 2) vector was excited. These precipitates are uniformly distributed in the magnesium matrix. Since the SAD pattern from the [1 1 2¯ 0] zone axis shows neither extra reflections nor diffuse scattering, it can be concluded that precipitates are fully coherent with the magnesium matrix. The precipitates are parallel to each other and perpendicular to the [0 0 0 1] direction. An elastic distortion, expected around precipitates with a value of around the 8%, is mostly adopted by a magnesium matrix [10]. Thus, an increase in the alloy hardness as the precipitates grow in the basal plane is expected. Fig. 6 shows the microhardness evolution during an isothermal treatment at 553 K. As can be seen, the alloy exhibits high hardness values under the as-deposited condition due to the solid solution reinforcement effect of alloying elements. As the thermal treatment progresses, the hardness increases due to the nucleation and growth of the titanium plates in the basal plane reaching a maximum after 30 h of annealing. Then, a slight decrease takes place due to coarsening of the titanium precipitates. The kinetic parameters of the process have been calculated to compare the precipitation reaction in this first peak with respect to the binary Mg–Ti alloy. The activation energies for the transformation can be determined by means of Kissinger’s analysis [16] as follows: Tp 2 E E ln (1) + ln ≈ RTp RK0 Φ
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Fig. 7. Kissinger plot of the first DSC precipitation peak.
tivation energies can be calculated. The values of E and K0 obtained from the Kissinger fit for the first are 160 kJ mol−1 and 10−13 s−1 , respectively, similar to those obtained in the binary system Mg–12%Ti [9]. Thus, there is no influence of the other alloying elements on the precipitation of pure titanium. On the other hand, the sample quenched at the end of the DSC scan shows coarser precipitates and extra double diffraction spots come out in the SAD pattern of the [1 1 2¯ 0] zone axis (Fig. 5b). The thickness of these precipitates in the [0 0 0 1] direction increases with respect to the initial precipitates (5–10 nm). Moreover, the precipitates are responsible for the appearance of Moir´e fringes when the (0 0 0 2) g vector is excited. In the case of the parallel Moir´e, the fringe spacing D of the structure precipitate-matrix is given by [17]: Fig. 5. (a–c) Bright field images and SADP from the [] zone axis showing the morphology of the titanium precipitates of the sample (a) quenched at first precipitation peak and (b) at the end of the DSC. Bright field image (c) and SADP (d) from the [0 0 1] zone axis of the microstructure of the sample quenched at the end of the DSC.
where E is the apparent activation energy, Tp the temperature of the maximum reaction rate, K0 the pre-exponential factor, R the universal gas constant and Φ the heating rate. The apparent activation energies for the transformation is obtained from the slope of the plot of ln (Tp 2 /Φ) versus 1/ Tp as shown in Fig. 7 for the case of the first peak where the temperature of the maximum is clearly distinguished. This figure shows a good linear correlation and the values of the apparent ac-
Fig. 6. Microhardness evolution during an isothermal treatment at 553 K.
1 1 1 = − D d1 d2
(2)
where d1 and d2 are the interplanar distance of the two superimposed crystal lattices. From Fig. 5b, the values for D can be measured obtaining an average value of 2.65 nm. Therefore, it is possible to calculate the interplanar distance of the precipitates taking as the interplanar distance between the (0 0 0 2) magnesium planes its theoretical value (0.260 nm). The value obtained for d2 is 0.236 and it is in concord with the theoretical interplanar distance between the (0 0 0 2) ␣-titanium planes. The epitaxial relationship of these titanium precipitates is expressed by [0 0 0 1]Ti ||[0 0 0 1] Mg and [1 1 2¯ 0]Ti ||[1 1 2¯ 0] Mg according with previous works [8–10]. It is worthy to note that the titanium precipitates present a hexagonal feature with their face parallel to the {1 1 2¯ 0} planes when the sample is observed in the [0 0 0 1] zone axis (Fig. 5c) and ellipsoidal with faces parallel to the traces of the {0 0 0 1} and {0 1 1¯ 1} planes when the sample is observed in the [1 1 2¯ 0] zone axis (Fig. 5b). During the precipitate coarsening the coherence between the titanium particles and the magnesium matrix is lost, provoking the appearance of misfit dislocation between the precipitates as it is observed in the weak beam image in Fig. 8.
G. Garc´es et al. / Journal of Alloys and Compounds 387 (2005) 115–120
Fig. 8. Weak beam image of the microstructure of the sample at the end of the DSC scan.
During the non-isothermal treatment, precipitation of the alloying elements takes place. The first two exothermal peaks are due to titanium precipitation, which takes place in two sequences as reported in binary Mg–Ti alloys. In the first peak, it can be supposed that titanium precipitates initially nucleate and grow to occupy a certain area in the (0 0 0 1) planes. The presence of small precipitates, which were observed when the (0 0 0 2) g vector was excited supports this argument. When the increase in misfit energy associated with small precipitate
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(basal plane) growth cannot be accomodated by the matrix, these precipitates grow along the [0 0 0 1] direction loosing their coherence with the matrix (second peak). Hitherto, it is assumed that this precipitates was incoherent or semicoherent with the matrix. However, the appearance of misfit dislocations generated in the precipitate interfaces has not been reported. The appearance of these misfit dislocations (Fig. 8) supports the assessment that these precipitates are incoherent. Not only the titanium precipitates are present at this stage but also spherical particles of about 20 nm decorating the grain boundaries are visible as shown in Fig. 9. The EDS results indicate that these particles are titanium and aluminium rich. The precipitation at grain boundaries of spherical Ti3 Al precipitates above 750 K, is consistent with the ternary equilibrium phase diagram of Mg–Al–Ti [18]. The isothermal section above 700 K shows that an alloy of similar composition as the alloy of this work would be composed of ␣Ti and the ␣2 phase (Ti3 Al DO19 ) in a magnesium matrix. This observation also agrees with recently studies of Wang et al. [19] where magnesium reinforced by in situ TiC particles was synthesized. Titanium, carbon and aluminium powders and Mg–9% Al–1% Zn (wt.%) alloy were used as the reinforced phase precursor and the matrix, respectively. During the composite processing not only TiC particles but also Ti3 Al precipitates was formed.
Fig. 9. Bright field image showing spherical precipitates at the grain boundaries. EDS spectra of the Mg matrix, spherical precipitates at grain boundaries and Ti precipitates within the grain.
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Finally, the continuous increase in the heat flow above 703 K of Fig. 3 is connected with the oxidation of the deposit [3,8]. The high concentration of defects, especially pores at the grain boundaries, induce rapid inward oxygen transport provoking the increase in the heat flow and mass gain in these PVD thick deposits [20,21].
4. Summary The thermal stability and the precipitation reaction in the PVD Mg–30% Ti–2% Al–0.9% Mn (wt.%) alloy have been studied. The microstructure of the alloy throughout the thickness consists of columnar grains in the deposit growth direction constituted by a solid solution of alloying elements in the magnesium matrix that provokes a strong contraction of the magnesium lattice parameters. During DSC experiments three exothermal transformation that correspond to precipitation phenomena in the magnesium matrix have been observed. The precipitation of titanium occurs like in the binary Mg–Ti alloy, in two steps. Furthermore, the presence of aluminium leads to the formation of the intermetallic phase Ti3 Al at the grain boundaries.
Acknowledgements The authors would like to thank Dr. S.B. Dodd, S. Morris and R.C. Piller of DERA (Qinetiq, actually), Farnborough, for providing the vapour deposited alloys along with useful information on processing conditions. We gratefully acknowledge the support of the Comisi´on Interministerial de Ciencia y Tecnolog´ıa (MAT 981620-CE) and Comunidad de Madrid (CAM 700-739).
References [1] S.B. Dodd, R.W. Gardiner, in: Proceedings of the Third International Magnesium Conference, The Institute of Metals, Manchester, 1996, p. 271. [2] S.B. Dodd, S. Morris, R.W. Gardiner, in: Proceedings of the Magnesium Alloys and their Applications, Frankfurt: WerkstoffInformationgesellschaft, 1998, p. 375. [3] S. Diplas, P. Tsakiropoulos, R.M.D. Brydson, Mater. Sci. Technol. 14 (1998) 689–698. [4] K.R. Baldwin, D.J. Bray, G.D. Howard, R.W. Gardiner, Mater. Sci. Technol. 12 (1996) 937–943. [5] S. Morris, S.B. Dodd, P.J. Hall, A.J. Mackinnom, L.E. Berlouis, J. Alloys Compd. 293–295 (1999) 458–462. [6] P. Adeva, S.B. Dodd, P. Morgan, F. Hehmann, P. Steinmetz, F. Sommer, Adv. Eng. Mater. 3 (2001) 147–152. [7] C.B. Baliga, P. Tsakiropoulos, S.B. Dodd, R.W. Gardiner, in: Proceedings of the Third International Magnesium Conference, The Institute of Metals, Manchester, 1996, p. 627. [8] K.E. Bagnall, P.G. Partridge, J.W. Steed, S.B. Dodd, R.W. Gardiner, in: Proceedings of the Third International Magnesium Conference, The Institute of Metals, Manchester, 1996, p. 299. [9] G. Garc´es, P. Adeva, Phil. Mag. A 82 (2002) 699–714. [10] G. Garc´es, P. P´erez, P. Adeva, Scripta Mater. 45 (2001) 1001–1007. [11] B.A. Movchan, A.V. Demchishin, Phys. Metal. Metallogr. 28 (1969) 653–660. [12] J.A. Thornton, Rev. Mater. Sci. 7 (1977) 239–260. [13] C.R.M. Grovenor, H.T.G. Hertzell, D.A. Smith, Acta Metall. 32 (1984) 773–781. [14] S. Diplas, P. Tsakiropoulos, R.M.D. Brydson, Mater. Sci. Technol. 15 (1999) 1359–1372. [15] J.P. Chu, C.H. Chung, P.Y. Lee, J.M. Rigsbee, J.Y. Wang, Metall. Mater. Trans. A 29A (1998) 647–658. [16] E.J. Mittemeijer, J. Mater. Sci. 27 (1992) 3977–3987. [17] D.B. Williams, C.B. Carter, Transmission Electron Microscopy, Plenum Press, New York, 1996. [18] G. Petzow, G. Effenberg, Ternary Alloys, vol. 6, ASM International, New York, Basel, 1993, p. 443. [19] H.Y. Wang, Q.C. Jiang, X.L. Li, J.G. Wang, Scripta Mater. 48 (2003) 1349–1354. [20] P. P´erez, G. Garc´es, P. Adeva, Oxid. Met. 58 (2002) 60. [21] P. P´erez, G. Garc´es, P. Adeva, Mater. Corr. 54 (2003) 145.