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Characterization of a laser-fabricated hypereutectic Al–Sc alloy bar
Available online at www.sciencedirect.com
ScienceDirect Scripta Materialia 87 (2014) 13–16 www.elsevier.com/locate/scriptamat
Characterization of a laser-fabricated hypereutectic Al–Sc alloy bar Paul A. Rometsch,a,⇑ Hao Zhong,a Kate M. Nairn,a Tom Jarvisa,b and Xinhua Wua,b a
b
Department of Materials Engineering, Monash University, Clayton, VIC 3800, Australia Monash Centre for Additive Manufacturing, 11 Normanby Road, Notting Hill, VIC 3168, Australia Received 16 April 2014; revised 27 May 2014; accepted 31 May 2014 Available online 12 June 2014
A 115 mm high bar was fabricated by laser metal deposition from Al–0.9 wt.% Sc powder. In the as-fabricated condition, the hardness decreased from 78 HV5 at the bottom to 34 HV5 at the top of the bar. At the top, the electrical conductivity was lowest, the amount of Sc in solution was highest and a subsequent ageing response was observed. It is concluded that for laser-fabrication of such alloys, the thermal history over the whole component must be understood and controlled. Ó 2014 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Laser deposition; Aluminum alloys; Nuclear magnetic resonance (NMR); Scandium
Although there is significant ongoing interest in the development of Al–Sc alloys due to their extraordinary properties, their use has been limited to several high-performance applications particularly in the aerospace and sporting goods industries because of the high price and uncertain supply of Sc [1–4]. Small Sc additions (typically <0.4 wt.%) to Al alloys can give significant: Grain refinement [2,5], Precipitation hardening from Al3Sc particles below 10 nm in size, reportedly capable of increasing the yield strength of pure Al from 15 to 199 MPa after alloying with 0.38 wt.% Sc, casting into a conventional ingot and ageing for 8 h at 288 °C [6], and Recrystallization prevention, which facilitates significant substructural hardening after deformation due to Al3Sc dispersoid particles being potent recrystallization and grain growth inhibitors [2,7,8]. As a result, Al–Mg–Sc alloys can achieve yield strengths in excess of 400 MPa after cold rolling and ageing [2]. Furthermore, exceptional combinations of strength, elongation, fracture toughness, fatigue resistance, electrical conductivity, corrosion resistance, weldability, superplasticity, creep resistance and/or elevated temperature stability can be achieved [1–4,7–12].
⇑ Corresponding
author. [email protected]
Tel.:
+61
399053655; e-mail:
paul.
With Al–Sc alloys, relatively fast solidification and cooling rates are desirable to avoid the formation of coarse Sc-containing particles and to keep as much Sc in solution as possible. Since the solvus already approaches the eutectic melting temperature below 0.4 wt.% Sc, a solution treatment is normally not able to place more Sc into solution and may even decrease the amount of Sc in solution, depending on how rapidly the alloy was solidified [2]. This was confirmed by nuclear magnetic resonance (NMR) experiments on an Al–0.49 wt.% Sc alloy, where 31.5% of all Sc atoms were still found to exist within Al3Sc precipitates even after solutionizing at 650 °C and water quenching [13]. The NMR method is particularly well suited for detecting minute amounts of Sc in Al and distinguishing between Sc in solution and Sc in precipitates [13]. Rapid solidification techniques such as melt spinning and spray forming have been used recently to place significantly more than 0.4 wt.% Sc into solution and thereby to increase the yield strength of Al–Mg–Sc alloys to 500 MPa or more after ageing, along with 10–15% elongations to fracture and very good notch toughness and fatigue properties [4,14–16]. Recently, similar property levels were achieved for an Al– 4.5Mg–0.66Sc–0.37Zr–0.51Mn alloy (wt.%) produced by selective laser melting (SLM) and subsequent artificial ageing for 4 h at 325 °C [17]. It has been claimed that such laser fabrication can result in solidification cooling rates well in excess of 1000 K/s [18,19], which would be ideal for hypereutectic Al–Sc alloys.
http://dx.doi.org/10.1016/j.scriptamat.2014.05.021 1359-6462/Ó 2014 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
14
P. A. Rometsch et al. / Scripta Materialia 87 (2014) 13–16
What has not been documented very well, however, is how the thermal history resulting from the layer-by-layer heat input in such laser additive manufacturing processes can change the microstructure and the properties along the build height of a precipitation hardening alloy. This work seeks to address these questions for a hypereutectic Al–Sc alloy by inference from mechanical testing, electrical conductivity, microscopy and NMR results. Gas atomized Al–Sc powder was fed into a 4 kW Trumpf 7040 blown powder laser metal deposition (LMD) machine. Over a 45-min period, a 15 mm wide 15 mm deep 115 mm high bar was fabricated layer-by-layer onto a 20 °C Al substrate in an argon atmosphere with a bidirectional scan strategy using a laser power of 800 W, a laser speed of 13 mm/s, a laser spot diameter of 2 mm and a powder layer thickness of 0.3 mm. The bar was then wire-cut into vertical slices, 1.3 mm thick, 15 mm wide and 115 mm high. This enabled the microstructure and properties to be evaluated along the build height from the first build layers at the bottom (near the Al substrate) to the final build layers at the top of the bar. Tensile testing was carried out on the flat samples using an Instron 4505/5500R screw-driven tensile machine at an extension rate of 1.2 mm/min and with sample gauge sections of 30 6 1.3 mm. Hardness testing was performed on a Vickers hardness tester using a 5 kg load. The electrical conductivity was measured at room temperature with a Foerster Sigmatest 2.068 eddy current device. Microstructural observations were made using a JEOL 7001F scanning electron microscope (SEM). Some samples were first aged in a salt bath at 325 °C, followed by water quenching. Nuclear magnetic resonance (NMR) spectroscopy was performed on both the gas atomized powder and on filings from the laser fabricated sample. The powder and/or filings (in the as-filed condition) were placed into a 4 mm diameter zirconia rotor and 45Sc spectra were measured on a Bruker Avance400 spectrometer operating at 9.4T. Samples were not spun, and a two-pulse echo sequence was used to minimize distortions from instrument deadtime. Spectra were referenced to ScCl3 (aq) at 0 ppm. Peak assignments and NMR methods were based on those used by Celotto and Bastow [13]. The gas atomized powder particle size ranged from 45 to 150 lm. The chemical compositions (in wt.%) of the powder and laser fabricated bar were determined by ICP-AES to be Al–0.90Sc–0.10Fe–0.07Si–0.03Mg and Al–0.87Sc–0.09Fe–0.07Si–0.03Mg, respectively.
This shows that the composition before and after laser fabrication was virtually unchanged. Furthermore, no compositional gradients were observed across the build. Hardness and electrical conductivity results along a vertical section of the bar are shown in Fig. 1 for both the as-fabricated condition and after ageing for up to eight days at 325 °C. In the as-fabricated condition, the hardness decreases from 78 HV5 at the bottom to 34 HV5 at the top of the bar. Upon ageing, there is no ageing response up to a height of about 70 mm. It is only in the top 30 mm of the bar that an ageing response is evident, with a maximum increase in hardness from 34 to 65 HV5 observed at the top of the bar after ageing for 2 h at 325 °C. Upon tensile testing the flat vertically oriented samples, fracture always occurred on the end of the gauge section that was 70 mm from the bottom and therefore close to the weakest point. The average tensile properties based on two samples were Rp0.2 = 97 ± 2 MPa, Rm = 135 ± 7 MPa, Au = 3.5 ± 0.7% and At = 5.6 ± 0.7% in the as-fabricated condition, and Rp0.2 = 79 ± 4 MPa, Rm = 116 ± 10 MPa, Au = 6.5 ± 0.5% and At = 10.0 ± 0.3% after ageing for 24 h at 325 °C (with errors representing maximum and minimum measured values). There is some variability in these results compared to the hardness results, and we attribute this to small differences in the exact fracture location. The electrical conductivity in the as-fabricated condition increased from 56%IACS at the bottom to 58%IACS at a height of 50 mm, and then decreased to 53.3%IACS at the top of the bar. Upon ageing, the maximum increase in conductivity was <1.3%IACS at the bottom of the bar and >6%IACS at the top of the bar. Since the electrical conductivity is known to decrease strongly with increasing solute in solution, these results show clearly that there is more Sc in solution at the top of the as-fabricated bar compared to further down the bar. This is supported by the hardness results, which indicate that there is insufficient Sc in solution to cause any appreciable amount of fine-scale Al3Sc hardening precipitates to form during ageing up to a height of 70 mm. The small increase in conductivity with increasing ageing time towards the bottom of the bar could be due to a decrease in the number density of very small Al3Sc precipitates as they coarsen [20]. At the top of the bar, the hardness increases more quickly than the conductivity within the first 12 min at 325 °C (by 19 HV5 vs. 1.9%IACS, respectively). The hardness then quickly reaches a relatively flat peak at 2–24 h, whereas the conductivity keeps increasing until
Figure 1. Variation in hardness and electrical conductivity from bottom to top of the bar in as-fabricated and artificially aged conditions. Typical error bars (variation from mean) for the hardness and electrical conductivity data are ±3 HV5 and ±0.1%IACS, respectively.
P. A. Rometsch et al. / Scripta Materialia 87 (2014) 13–16
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Figure 2. Backscattered SEM images at two different magnifications showing how the as-fabricated cross-sectional microstructure changes with increasing distance up from the bottom of the bar: (a, b) 10 mm, (c, d) 70 mm, and (e, f) 110 mm. All scale bars represent 1 lm.
the material is slightly over-aged after eight days at 325 °C. The rapid age hardening at the beginning is due to the formation of a high number density of nanometer-sized Al3Sc hardening precipitates. The initial rate of increase in the conductivity is a little slower because the positive effect of decreasing Sc in solution is partly counteracted by the interfacial scattering of electrons because of the high number density of nano-scale Al3Sc precipitates [20]. Upon longer ageing, the conductivity keeps increasing and the hardness drops as the amount of Sc in solution decreases and the Al3Sc precipitates coarsen and decrease in number. The SEM images in Fig. 2 show fine micron-sized grains and precipitates that were found by EDXS to contain Sc and/or Fe. The particles are resolved as bright white phases in the backscattered electron images because of the relatively high atomic numbers of Sc and Fe compared to Al. They are clearly coarser at the top of the bar than at the bottom, suggesting that there is a progression from more high temperature exposure at the top to less high temperature exposure at the bottom of the bar. Interestingly, this resulted in a significant ageing response during subsequent ageing at the top of the bar but not further down the bar (Fig. 1). The reason for this is that the heat from building the upper layers is conducted down through the bar and into the Al substrate. This effectively induces a temperature gradient down through the bar such that the average temperature decreases with increasing distance from the laser melt pool. As a result, it is expected that a significant amount of ageing is already occurring in the lower layers of the bar during prolonged thermal cycling (over a range of, say, 150–400 °C) while the upper layers are still being built, and that the degree of this ageing decreases with increasing distance up the bar. There is some evidence
for this in Fig. 2(b) in the form of the small white spots that are believed to be 20–70 nm Al3Sc precipitates, though the smaller hardening precipitates are not resolvable at this level of SEM. By contrast, there were no such fine precipitates at the top of the bar. The bottom of the bar therefore achieved a higher hardness for three main reasons: (i) it experienced faster cooling rates due to its proximity to the unheated Al substrate, (ii) it experienced less elevated temperature exposures than further up the bar, thereby resulting in a reduced loss of Sc to coarse non-hardening precipitates, and (iii) it was effectively aged more than the middle and top of the bar during laser fabrication. This explanation is broadly consistent with the temperature–time measurement and modelling trends on small Ti–6Al–4V samples fabricated on a similar blown powder LMD machine [18], although the actual temperatures will differ. The above explanations are confirmed by our NMR results. To evaluate the fractions of Sc in solution and in Al3Sc, the area ratios of the two NMR peaks were taken. Attempts to deconvolute the spectra by peak fitting were not successful, so the areas were evaluated manually, by cutting up printouts of the spectra and weighing the peaks. This method slightly underestimates the areas of the smaller peaks in the spectra. Firstly, NMR of the rapidly solidified gas atomized powder revealed that 91% of the available Sc (i.e. 0.79 wt.%) was in a supersaturated solid solution. Secondly, in the as-fabricated condition, there was almost triple the Sc in solution at the top of the bar compared to at the bottom of the bar (Fig. 3). Furthermore, there were similarly low levels of Sc in solution at 10 mm and 70 mm distances up the bar. This confirms that so much ageing has already occurred during fabrication at both 10 mm and 70 mm distances up the bar that there is not enough Sc remaining in solution
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P. A. Rometsch et al. / Scripta Materialia 87 (2014) 13–16
Figure 3. 45Sc NMR spectra and amounts of Sc in solid solution shown for: (a) gas atomized powder, (b) filings from 10 mm up the asfabricated bar, (c) filings from 70 mm up the as-fabricated bar, (d) filings from 110 mm up the as-fabricated bar, and (e) filings from 110 mm up the bar after ageing for 24 h at 325 °C. The two peaks are assigned to Sc in solution (1700 ppm) and in Al3Sc precipitates (1050 ppm).
to cause any appreciable subsequent ageing response in terms of hardness. The 1.5%IACS lower electrical conductivity at the 10 mm location compared to the 70 mm location is probably because of more small Sc-containing clusters/precipitates at the 10 mm location due to the lower effective ageing temperatures (from thermal cycling) there. The hump in conductivity between top and bottom therefore results from two competing effects during laser fabrication: (i) more Sc in solution from higher temperature thermal cycling at the top of the bar, and (ii) more small precipitates from lower temperature ageing further down the bar. The third finding from the NMR work is that the amount of Sc in solution at the top of the bar does indeed decrease after subsequent ageing for 24 h at 325 °C (Fig. 3). This explains why a significant ageing response is observed at the top of the bar in terms of both hardness and electrical conductivity. Based on the areas under the spectra in Fig. 3(d) and (e), the amount of Sc in solution has decreased from 0.11 to 0.08 wt.% after ageing. This ageing treatment almost doubled the hardness; however, even after this ageing treatment the amount of Sc in solution is still greater than the amount remaining in solution at 10 and 70 mm distances up the bar in the as-fabricated condition. Nevertheless, in all cases (apart from the gas atomized powder), the amount of Sc contained in Al3Sc precipitates is significantly greater than the amount of Sc in solution. These results confirm that a significant amount of expensive Sc was lost to coarse, non-hardening precipitates. The significant variations in microstructure and properties reported here along the LMD build height are at odds with the recently reported high degrees of property uniformity across Al–Mg–Sc–Zr parts built by SLM [17,19]. However, compared to a typical SLM process, the current LMD process achieved a 150 times slower laser speed, a 20 times larger laser spot diameter, a 2.2 times higher laser power, a 7.5 times greater powder layer thickness and consequently more than twice the energy input per unit volume of melted powder (i.e. an estimated 103 J/mm3 for LMD vs. 47 J/mm3 for SLM). It is therefore not unreasonable to expect the SLM process to result in less microstructure and property variation along the build height. Overall, this work demonstrates that while a significant amount of hardening from Al3Sc precipitation is possible in laser fabricated Al–Sc alloys either with or
without subsequent ageing, care must be taken if good combinations of properties are to be obtained. In particular, it was found that the hardness, electrical conductivity, Sc content in solution and microstructure vary significantly from the bottom to the top of a 115 mm high laser metal deposited bar due to a complex progression of temperature–time thermal cycling exposures along the build height. If greater uniformity and better properties are to be achieved, then it is crucial to understand and control the thermal history over the whole component. A first step would be to avoid the formation of coarse non-hardening Sc-containing precipitates by adjusting the processing parameters to reduce the amount of elevated temperature thermal cycling over successive build layers. We are grateful for access to the facilities at the Monash Centre for Additive Manufacturing (MCAM) and the Monash Centre for Electron Microscopy (MCEM). The gas atomized powder was provided under the European Accelerated Metallurgy Project framework. We thank Dr. Tim Bastow for helpful discussions and assistance with the NMR work at CSIRO in Clayton, Melbourne. [1] L.S. Toropova, D.G. Eskin, M.L. Kharakterova, T.V. Dobatkina, Advanced Aluminum Alloys Containing Scandium – Structure and Properties, Gordon and Breach Science Publishers, Amsterdam, 1998. [2] J. Røyset, N. Ryum, Int. Mater. Rev. 50 (2005) 19. [3] F. Palm, International Patent: WO 03/052154 (2003). [4] F. Palm, P. Vermeer, W. von Bestenbostel, D. Isheim, R. Schneider, in: Proc. of 11th Intern. Conf. on Aluminium Alloys (ICAA11), Aachen, Germany, DGM and WileyVCH, 2008. [5] Y.C. Ye, L.J. He, P.J. Li, Trans. Non-ferr. Met. Soc. China 20 (2010) 465. [6] L.A. Willey, US Patent: 3,619,181 (1971). [7] V.I. Elagin, Met. Sci. Heat Treat. 49 (2007) 427. [8] V.V. Zakharov, T.D. Rostova, Met. Sci. Heat Treat. 49 (2007) 435. [9] E.A. Marquis, D.N. Seidman, D.C. Dunand, Acta Mater. 51 (2003) 4751. [10] D.N. Seidman, E.A. Marquis, D.C. Dunand, Acta Mater. 50 (2002) 4021. [11] M.E. van Dalen, D.N. Seidman, D.C. Dunand, Acta Mater. 56 (2008) 4369. [12] K.E. Knipling, R.A. Karnesky, C.P. Lee, D.C. Dunand, D.N. Seidman, Acta Mater. 58 (2010) 5184. [13] S. Celotto, T.J. Bastow, Philos. Mag. A 80 (2000) 1111. [14] T. Herding, O. Kebler, H.-W. Zoch, Materialwiss. Werkstofftech. 38 (2007) 855. [15] F. Palm, R. Leuschner, T. Schubert, B. Kieback, in: Proc. of Powder Metall. World Congr. (World PM2010), Florence, Italy, European Powder Metallurgy Association, 2010. [16] A.B. Pandey, European Patent: EP 2251447 (2010). [17] K. Schmidtke, F. Palm, A. Hawkins, C. Emmelmann, Phys. Procedia 12 (2011) 369. [18] L. Qian, J. Mei, J. Liang, X. Wu, Mater. Sci. Technol. 21 (2005) 1. [19] F. Palm, K. Schmidtke, in: Proc. of 9th Intern. Conf. on Trends in Welding Research, Chicago, IL, ASM International, 2012. [20] F. Fazeli, C.W. Sinclair, T. Bastow, Metall. Mater. Trans. A 39 (2008) 2297.
ScienceDirect Scripta Materialia 87 (2014) 13–16 www.elsevier.com/locate/scriptamat
Characterization of a laser-fabricated hypereutectic Al–Sc alloy bar Paul A. Rometsch,a,⇑ Hao Zhong,a Kate M. Nairn,a Tom Jarvisa,b and Xinhua Wua,b a
b
Department of Materials Engineering, Monash University, Clayton, VIC 3800, Australia Monash Centre for Additive Manufacturing, 11 Normanby Road, Notting Hill, VIC 3168, Australia Received 16 April 2014; revised 27 May 2014; accepted 31 May 2014 Available online 12 June 2014
A 115 mm high bar was fabricated by laser metal deposition from Al–0.9 wt.% Sc powder. In the as-fabricated condition, the hardness decreased from 78 HV5 at the bottom to 34 HV5 at the top of the bar. At the top, the electrical conductivity was lowest, the amount of Sc in solution was highest and a subsequent ageing response was observed. It is concluded that for laser-fabrication of such alloys, the thermal history over the whole component must be understood and controlled. Ó 2014 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Laser deposition; Aluminum alloys; Nuclear magnetic resonance (NMR); Scandium
Although there is significant ongoing interest in the development of Al–Sc alloys due to their extraordinary properties, their use has been limited to several high-performance applications particularly in the aerospace and sporting goods industries because of the high price and uncertain supply of Sc [1–4]. Small Sc additions (typically <0.4 wt.%) to Al alloys can give significant: Grain refinement [2,5], Precipitation hardening from Al3Sc particles below 10 nm in size, reportedly capable of increasing the yield strength of pure Al from 15 to 199 MPa after alloying with 0.38 wt.% Sc, casting into a conventional ingot and ageing for 8 h at 288 °C [6], and Recrystallization prevention, which facilitates significant substructural hardening after deformation due to Al3Sc dispersoid particles being potent recrystallization and grain growth inhibitors [2,7,8]. As a result, Al–Mg–Sc alloys can achieve yield strengths in excess of 400 MPa after cold rolling and ageing [2]. Furthermore, exceptional combinations of strength, elongation, fracture toughness, fatigue resistance, electrical conductivity, corrosion resistance, weldability, superplasticity, creep resistance and/or elevated temperature stability can be achieved [1–4,7–12].
⇑ Corresponding
author. [email protected]
Tel.:
+61
399053655; e-mail:
paul.
With Al–Sc alloys, relatively fast solidification and cooling rates are desirable to avoid the formation of coarse Sc-containing particles and to keep as much Sc in solution as possible. Since the solvus already approaches the eutectic melting temperature below 0.4 wt.% Sc, a solution treatment is normally not able to place more Sc into solution and may even decrease the amount of Sc in solution, depending on how rapidly the alloy was solidified [2]. This was confirmed by nuclear magnetic resonance (NMR) experiments on an Al–0.49 wt.% Sc alloy, where 31.5% of all Sc atoms were still found to exist within Al3Sc precipitates even after solutionizing at 650 °C and water quenching [13]. The NMR method is particularly well suited for detecting minute amounts of Sc in Al and distinguishing between Sc in solution and Sc in precipitates [13]. Rapid solidification techniques such as melt spinning and spray forming have been used recently to place significantly more than 0.4 wt.% Sc into solution and thereby to increase the yield strength of Al–Mg–Sc alloys to 500 MPa or more after ageing, along with 10–15% elongations to fracture and very good notch toughness and fatigue properties [4,14–16]. Recently, similar property levels were achieved for an Al– 4.5Mg–0.66Sc–0.37Zr–0.51Mn alloy (wt.%) produced by selective laser melting (SLM) and subsequent artificial ageing for 4 h at 325 °C [17]. It has been claimed that such laser fabrication can result in solidification cooling rates well in excess of 1000 K/s [18,19], which would be ideal for hypereutectic Al–Sc alloys.
http://dx.doi.org/10.1016/j.scriptamat.2014.05.021 1359-6462/Ó 2014 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
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P. A. Rometsch et al. / Scripta Materialia 87 (2014) 13–16
What has not been documented very well, however, is how the thermal history resulting from the layer-by-layer heat input in such laser additive manufacturing processes can change the microstructure and the properties along the build height of a precipitation hardening alloy. This work seeks to address these questions for a hypereutectic Al–Sc alloy by inference from mechanical testing, electrical conductivity, microscopy and NMR results. Gas atomized Al–Sc powder was fed into a 4 kW Trumpf 7040 blown powder laser metal deposition (LMD) machine. Over a 45-min period, a 15 mm wide 15 mm deep 115 mm high bar was fabricated layer-by-layer onto a 20 °C Al substrate in an argon atmosphere with a bidirectional scan strategy using a laser power of 800 W, a laser speed of 13 mm/s, a laser spot diameter of 2 mm and a powder layer thickness of 0.3 mm. The bar was then wire-cut into vertical slices, 1.3 mm thick, 15 mm wide and 115 mm high. This enabled the microstructure and properties to be evaluated along the build height from the first build layers at the bottom (near the Al substrate) to the final build layers at the top of the bar. Tensile testing was carried out on the flat samples using an Instron 4505/5500R screw-driven tensile machine at an extension rate of 1.2 mm/min and with sample gauge sections of 30 6 1.3 mm. Hardness testing was performed on a Vickers hardness tester using a 5 kg load. The electrical conductivity was measured at room temperature with a Foerster Sigmatest 2.068 eddy current device. Microstructural observations were made using a JEOL 7001F scanning electron microscope (SEM). Some samples were first aged in a salt bath at 325 °C, followed by water quenching. Nuclear magnetic resonance (NMR) spectroscopy was performed on both the gas atomized powder and on filings from the laser fabricated sample. The powder and/or filings (in the as-filed condition) were placed into a 4 mm diameter zirconia rotor and 45Sc spectra were measured on a Bruker Avance400 spectrometer operating at 9.4T. Samples were not spun, and a two-pulse echo sequence was used to minimize distortions from instrument deadtime. Spectra were referenced to ScCl3 (aq) at 0 ppm. Peak assignments and NMR methods were based on those used by Celotto and Bastow [13]. The gas atomized powder particle size ranged from 45 to 150 lm. The chemical compositions (in wt.%) of the powder and laser fabricated bar were determined by ICP-AES to be Al–0.90Sc–0.10Fe–0.07Si–0.03Mg and Al–0.87Sc–0.09Fe–0.07Si–0.03Mg, respectively.
This shows that the composition before and after laser fabrication was virtually unchanged. Furthermore, no compositional gradients were observed across the build. Hardness and electrical conductivity results along a vertical section of the bar are shown in Fig. 1 for both the as-fabricated condition and after ageing for up to eight days at 325 °C. In the as-fabricated condition, the hardness decreases from 78 HV5 at the bottom to 34 HV5 at the top of the bar. Upon ageing, there is no ageing response up to a height of about 70 mm. It is only in the top 30 mm of the bar that an ageing response is evident, with a maximum increase in hardness from 34 to 65 HV5 observed at the top of the bar after ageing for 2 h at 325 °C. Upon tensile testing the flat vertically oriented samples, fracture always occurred on the end of the gauge section that was 70 mm from the bottom and therefore close to the weakest point. The average tensile properties based on two samples were Rp0.2 = 97 ± 2 MPa, Rm = 135 ± 7 MPa, Au = 3.5 ± 0.7% and At = 5.6 ± 0.7% in the as-fabricated condition, and Rp0.2 = 79 ± 4 MPa, Rm = 116 ± 10 MPa, Au = 6.5 ± 0.5% and At = 10.0 ± 0.3% after ageing for 24 h at 325 °C (with errors representing maximum and minimum measured values). There is some variability in these results compared to the hardness results, and we attribute this to small differences in the exact fracture location. The electrical conductivity in the as-fabricated condition increased from 56%IACS at the bottom to 58%IACS at a height of 50 mm, and then decreased to 53.3%IACS at the top of the bar. Upon ageing, the maximum increase in conductivity was <1.3%IACS at the bottom of the bar and >6%IACS at the top of the bar. Since the electrical conductivity is known to decrease strongly with increasing solute in solution, these results show clearly that there is more Sc in solution at the top of the as-fabricated bar compared to further down the bar. This is supported by the hardness results, which indicate that there is insufficient Sc in solution to cause any appreciable amount of fine-scale Al3Sc hardening precipitates to form during ageing up to a height of 70 mm. The small increase in conductivity with increasing ageing time towards the bottom of the bar could be due to a decrease in the number density of very small Al3Sc precipitates as they coarsen [20]. At the top of the bar, the hardness increases more quickly than the conductivity within the first 12 min at 325 °C (by 19 HV5 vs. 1.9%IACS, respectively). The hardness then quickly reaches a relatively flat peak at 2–24 h, whereas the conductivity keeps increasing until
Figure 1. Variation in hardness and electrical conductivity from bottom to top of the bar in as-fabricated and artificially aged conditions. Typical error bars (variation from mean) for the hardness and electrical conductivity data are ±3 HV5 and ±0.1%IACS, respectively.
P. A. Rometsch et al. / Scripta Materialia 87 (2014) 13–16
15
Figure 2. Backscattered SEM images at two different magnifications showing how the as-fabricated cross-sectional microstructure changes with increasing distance up from the bottom of the bar: (a, b) 10 mm, (c, d) 70 mm, and (e, f) 110 mm. All scale bars represent 1 lm.
the material is slightly over-aged after eight days at 325 °C. The rapid age hardening at the beginning is due to the formation of a high number density of nanometer-sized Al3Sc hardening precipitates. The initial rate of increase in the conductivity is a little slower because the positive effect of decreasing Sc in solution is partly counteracted by the interfacial scattering of electrons because of the high number density of nano-scale Al3Sc precipitates [20]. Upon longer ageing, the conductivity keeps increasing and the hardness drops as the amount of Sc in solution decreases and the Al3Sc precipitates coarsen and decrease in number. The SEM images in Fig. 2 show fine micron-sized grains and precipitates that were found by EDXS to contain Sc and/or Fe. The particles are resolved as bright white phases in the backscattered electron images because of the relatively high atomic numbers of Sc and Fe compared to Al. They are clearly coarser at the top of the bar than at the bottom, suggesting that there is a progression from more high temperature exposure at the top to less high temperature exposure at the bottom of the bar. Interestingly, this resulted in a significant ageing response during subsequent ageing at the top of the bar but not further down the bar (Fig. 1). The reason for this is that the heat from building the upper layers is conducted down through the bar and into the Al substrate. This effectively induces a temperature gradient down through the bar such that the average temperature decreases with increasing distance from the laser melt pool. As a result, it is expected that a significant amount of ageing is already occurring in the lower layers of the bar during prolonged thermal cycling (over a range of, say, 150–400 °C) while the upper layers are still being built, and that the degree of this ageing decreases with increasing distance up the bar. There is some evidence
for this in Fig. 2(b) in the form of the small white spots that are believed to be 20–70 nm Al3Sc precipitates, though the smaller hardening precipitates are not resolvable at this level of SEM. By contrast, there were no such fine precipitates at the top of the bar. The bottom of the bar therefore achieved a higher hardness for three main reasons: (i) it experienced faster cooling rates due to its proximity to the unheated Al substrate, (ii) it experienced less elevated temperature exposures than further up the bar, thereby resulting in a reduced loss of Sc to coarse non-hardening precipitates, and (iii) it was effectively aged more than the middle and top of the bar during laser fabrication. This explanation is broadly consistent with the temperature–time measurement and modelling trends on small Ti–6Al–4V samples fabricated on a similar blown powder LMD machine [18], although the actual temperatures will differ. The above explanations are confirmed by our NMR results. To evaluate the fractions of Sc in solution and in Al3Sc, the area ratios of the two NMR peaks were taken. Attempts to deconvolute the spectra by peak fitting were not successful, so the areas were evaluated manually, by cutting up printouts of the spectra and weighing the peaks. This method slightly underestimates the areas of the smaller peaks in the spectra. Firstly, NMR of the rapidly solidified gas atomized powder revealed that 91% of the available Sc (i.e. 0.79 wt.%) was in a supersaturated solid solution. Secondly, in the as-fabricated condition, there was almost triple the Sc in solution at the top of the bar compared to at the bottom of the bar (Fig. 3). Furthermore, there were similarly low levels of Sc in solution at 10 mm and 70 mm distances up the bar. This confirms that so much ageing has already occurred during fabrication at both 10 mm and 70 mm distances up the bar that there is not enough Sc remaining in solution
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Figure 3. 45Sc NMR spectra and amounts of Sc in solid solution shown for: (a) gas atomized powder, (b) filings from 10 mm up the asfabricated bar, (c) filings from 70 mm up the as-fabricated bar, (d) filings from 110 mm up the as-fabricated bar, and (e) filings from 110 mm up the bar after ageing for 24 h at 325 °C. The two peaks are assigned to Sc in solution (1700 ppm) and in Al3Sc precipitates (1050 ppm).
to cause any appreciable subsequent ageing response in terms of hardness. The 1.5%IACS lower electrical conductivity at the 10 mm location compared to the 70 mm location is probably because of more small Sc-containing clusters/precipitates at the 10 mm location due to the lower effective ageing temperatures (from thermal cycling) there. The hump in conductivity between top and bottom therefore results from two competing effects during laser fabrication: (i) more Sc in solution from higher temperature thermal cycling at the top of the bar, and (ii) more small precipitates from lower temperature ageing further down the bar. The third finding from the NMR work is that the amount of Sc in solution at the top of the bar does indeed decrease after subsequent ageing for 24 h at 325 °C (Fig. 3). This explains why a significant ageing response is observed at the top of the bar in terms of both hardness and electrical conductivity. Based on the areas under the spectra in Fig. 3(d) and (e), the amount of Sc in solution has decreased from 0.11 to 0.08 wt.% after ageing. This ageing treatment almost doubled the hardness; however, even after this ageing treatment the amount of Sc in solution is still greater than the amount remaining in solution at 10 and 70 mm distances up the bar in the as-fabricated condition. Nevertheless, in all cases (apart from the gas atomized powder), the amount of Sc contained in Al3Sc precipitates is significantly greater than the amount of Sc in solution. These results confirm that a significant amount of expensive Sc was lost to coarse, non-hardening precipitates. The significant variations in microstructure and properties reported here along the LMD build height are at odds with the recently reported high degrees of property uniformity across Al–Mg–Sc–Zr parts built by SLM [17,19]. However, compared to a typical SLM process, the current LMD process achieved a 150 times slower laser speed, a 20 times larger laser spot diameter, a 2.2 times higher laser power, a 7.5 times greater powder layer thickness and consequently more than twice the energy input per unit volume of melted powder (i.e. an estimated 103 J/mm3 for LMD vs. 47 J/mm3 for SLM). It is therefore not unreasonable to expect the SLM process to result in less microstructure and property variation along the build height. Overall, this work demonstrates that while a significant amount of hardening from Al3Sc precipitation is possible in laser fabricated Al–Sc alloys either with or
without subsequent ageing, care must be taken if good combinations of properties are to be obtained. In particular, it was found that the hardness, electrical conductivity, Sc content in solution and microstructure vary significantly from the bottom to the top of a 115 mm high laser metal deposited bar due to a complex progression of temperature–time thermal cycling exposures along the build height. If greater uniformity and better properties are to be achieved, then it is crucial to understand and control the thermal history over the whole component. A first step would be to avoid the formation of coarse non-hardening Sc-containing precipitates by adjusting the processing parameters to reduce the amount of elevated temperature thermal cycling over successive build layers. We are grateful for access to the facilities at the Monash Centre for Additive Manufacturing (MCAM) and the Monash Centre for Electron Microscopy (MCEM). The gas atomized powder was provided under the European Accelerated Metallurgy Project framework. We thank Dr. Tim Bastow for helpful discussions and assistance with the NMR work at CSIRO in Clayton, Melbourne. [1] L.S. Toropova, D.G. Eskin, M.L. Kharakterova, T.V. Dobatkina, Advanced Aluminum Alloys Containing Scandium – Structure and Properties, Gordon and Breach Science Publishers, Amsterdam, 1998. [2] J. Røyset, N. Ryum, Int. Mater. Rev. 50 (2005) 19. [3] F. Palm, International Patent: WO 03/052154 (2003). [4] F. Palm, P. Vermeer, W. von Bestenbostel, D. Isheim, R. Schneider, in: Proc. of 11th Intern. Conf. on Aluminium Alloys (ICAA11), Aachen, Germany, DGM and WileyVCH, 2008. [5] Y.C. Ye, L.J. He, P.J. Li, Trans. Non-ferr. Met. Soc. China 20 (2010) 465. [6] L.A. Willey, US Patent: 3,619,181 (1971). [7] V.I. Elagin, Met. Sci. Heat Treat. 49 (2007) 427. [8] V.V. Zakharov, T.D. Rostova, Met. Sci. Heat Treat. 49 (2007) 435. [9] E.A. Marquis, D.N. Seidman, D.C. Dunand, Acta Mater. 51 (2003) 4751. [10] D.N. Seidman, E.A. Marquis, D.C. Dunand, Acta Mater. 50 (2002) 4021. [11] M.E. van Dalen, D.N. Seidman, D.C. Dunand, Acta Mater. 56 (2008) 4369. [12] K.E. Knipling, R.A. Karnesky, C.P. Lee, D.C. Dunand, D.N. Seidman, Acta Mater. 58 (2010) 5184. [13] S. Celotto, T.J. Bastow, Philos. Mag. A 80 (2000) 1111. [14] T. Herding, O. Kebler, H.-W. Zoch, Materialwiss. Werkstofftech. 38 (2007) 855. [15] F. Palm, R. Leuschner, T. Schubert, B. Kieback, in: Proc. of Powder Metall. World Congr. (World PM2010), Florence, Italy, European Powder Metallurgy Association, 2010. [16] A.B. Pandey, European Patent: EP 2251447 (2010). [17] K. Schmidtke, F. Palm, A. Hawkins, C. Emmelmann, Phys. Procedia 12 (2011) 369. [18] L. Qian, J. Mei, J. Liang, X. Wu, Mater. Sci. Technol. 21 (2005) 1. [19] F. Palm, K. Schmidtke, in: Proc. of 9th Intern. Conf. on Trends in Welding Research, Chicago, IL, ASM International, 2012. [20] F. Fazeli, C.W. Sinclair, T. Bastow, Metall. Mater. Trans. A 39 (2008) 2297.