of WO3and CoO2as reported after sliding at high temperatures (Ref 4). On the plate specimens only very mild wear is observed, while for the counter ball specimens an increase in apparent contact area is observed for all the different roughness. The increase in apparent contact area after 720 cycles is greatest for the specimens with lowest roughness while after sliding 180 cycles the difference is negligible between the different samples.
For investigating the mechanical properties of the very top surface of the coating compression testing on micro pillars were done. Pillars were milled out by FIB and indented with a flat tip indenter. The pillars showed and elastic response to indentation until a critical load was applied and failure of the pillars occurred. Force versus displacement curves indicate
that the properties of the coating were fairly consistent in the elastic region, but exhibited different critical loads as seen by the four different loading to failure curves in Figure 5. SEM investigation of the pillars after compression testing indicated that failure occurred in the binder phase between the carbides as would be expected.
The difference in critical load may be a result of non-homogenous distribution of carbides and carbide size in the pillars as well as the presence of pores inside the pillars. Highest load before failure was at 10071 mN while the weakest pillar had a critical load of 6536 mN (Figure 4).
Figure 4: Load vs. displacement for compression of
micropillar.
The failure modes of the weakest pillar was deformation and sliding of the top surface while the strongest pillar showed a sudden catastrophic failure as seen in Fig. 5.
Figure5: SEM image of pillar failure after compression.
Summary and Conclusion
Reducing the surface roughness increases CoF for ball on plate experiments. Only mild polishing wear is observed in the wear tracks of the rough specimens similar to wear tracks on real gate components. The presence of a cracked tribofilm is not visible on real components as observed after ball on plate testing. Mechanical testing by compression of micro pillars is feasible for investigating the properties of the coating on selective areas.
References
1.N.-E. ISO, “NS-EN ISO 10423, Petroleum and
Natural Gas Industries, Drilling and Production
Equipment, Wellhead and Christmas Tree
Equipment," ed, 2009.
2.N.-E. ISO, "NS-EN ISO 13628-4, Petroleum and
Natural Gas Industries-Design and Operations of
Subsea Production Systems, Part 4: Subsea Wellhead
and Tree Equipment ", ed, 2010.
3.G. Bolelli, V. Cannillo, L. Lusvarghi, and T.
Manfredini, Wear Behaviour of Thermally Sprayed
Ceramic Oxide Coatings, Wear, vol. 261, pp. 1298-
1315, 2006.
4.G.M. Balamurugan, M. Duraiselvam, and V.
Anandakrishnan, Comparison of High Temperature
Wear Behaviour of Plasma Sprayed WC-Co Coated
and Hard Chromium Plated AISI 304 Austenitic
Stainless Steel, Materials and Design, vol. 35, pp.
640-646, 2012.
Nucleation and Growth Transformations in
Vacuum Plasma Sprayed Ti-6Al-4V Alloy
*H.R. Salimijazi, Z.A. Mousavi, M.A. Golozar
Materials Engineering Department, Isfahan University of Technology, 84156-83111, Isfahan, Iran,
*E-mail:jazi@mie.utoronto.ca
J. Mostaghimi, T, Coyle
Centre for Advanced Coating Technologies, University of Toronto, Toronto, Canada
Abstract
Because of the nature of the plasma spray processes, physical and mechanical properties of the vacuum plasma sprayed structures of Ti-6Al-4V alloys are completely different compared to conventionally manufactured alloy. In order to reach to desirable mechanical and physical properties, vacancy and internal defects must be reduced, splats boundaries must be eliminated, and the optimal phase compositions should be obtained through the post deposition heat treatment. In order to have appropriate heat treatment processes, it will be needed to study the kinetic behavior of the as-sprayed microstructure at elevated temperatures. In the current study, the kinetic of solid transformations in Ti-6Al-4V alloys produced by vacuum plasma spraying process was studied based on Johnson-Mehl-Avrami (JMA) theory. In the kinetic behavior of this alloy, the dependency and lack of equality of the transformation rate constant with temperature caused an irregularity at 900oC. This irregularity showed deference between transformation mechanism above and below 900oC. At lower temperature (<900oC) curves constant gradient showed lack of change in the transformation mechanism including homogeneous nucleation and grown of α-phase. At higher temperature (>900oC) the gradient change indicated change in the transformation mechanism (first mechanism was formation of α-phase grain boundary and second mechanism was α-plate nucleation and grown from the grain boundaries). The value of the transformation rate constant in the kinetic study of the as-sprayed Ti-6Al-4V alloy was much higher than that produced from casting method. By using the results obtained from kinetics of βαβphase transformation at different constant temperatures, TTT diagram for the as-sprayed Ti–6Al–4V alloy was developed
Introduction
Titanium and titanium alloys are finding widespread applications in aerospace industries for more than three decades due to their desirable combination of properties such as high specific strength to weight ratio, good corrosion resistance in many environments and high fatigue strength (Ref 1-3). Titanium is an allotropic element. At room temperature, titanium has a hexagonal close-packed (hcp) crystal structure, which is referred to as “alpha” phase. This structure transforms to a body-centered cubic (bcc) crystal structure, so called “beta” phase, at 883°C. Titanium alloys classified into three main groups: α alloys, α+β alloys, and β alloys (Ref 4). α+β alloys contain a mixture of α and β phases at room temperature. The α-β alloys have the greatest commercial importance with one composition, Ti-6Al-4V, making up more than half the sales of titanium alloys both in Europe and the United States (Ref 5). Manipulating the morphology of the two allotropic phases of α and β with various heat treatment cycles in Ti-6Al-4V alloy yields excellent combination of strength and toughness along with excellent corrosion resistance. Under equilibrium conditions, the transformation of the alloy from α+β to β structure occurs at β transition temperature, which is about 980oC (Ref 6).
Ti-6Al-4V components are conventionally manufactured by casting, forging, and powder metallurgy processes. The high reactivity of titanium in the molten state requires specific casting technology. The low thermal conductivity of titanium leads to the formation of a coarse-grained microstructure, gas absorption, and segregation, which results in a reduction in mechanical properties. Ti-6Al-4V castings are about two to three times the cost of super alloy castings because of the required process conditions. The difficulty of forging complex-shaped structures has limited the use of wrought Ti-6Al-4V in such applications. Powder metallurgy is used to produce near-net-shape components. This process produces a structure with less than full density which generally has lower ductility, toughness, and fatigue strength than wrought materials (Ref 5).
Vacuum plasma spray forming (VPSF) is a new technology for manufacturing mechanical components from metals, ceramics, and composites that cannot easily be manufactured using conventional methods. Near-net-shape manufacturing by VPSF combines the processes of melting, rapid solidification, consolidation, and welding in a single step without facing the obstacles encountered in casting, forging, powder metallurgy, and welding (Ref 7). In an ideal plasma spray process, powders are fed into the high-velocity, high-temperature gas jet, generated by a direct-current DC arc in a plasma gun, using a feeder system. Depending on the powder size
Thermal Spray 2013—Innovative Coating Solutions for the Global Economy Proceedings of the International Thermal Spray Conference
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distribution and plasma spraying process parameters, powders are completely or partially melted and accelerated toward the substrate or mandrel mold. Finally, the droplets solidify after impacting with a very high cooling rate ~–°C/s, forming lamellae splats. Coatings are formed by building up layers of individual splats (Ref 8, 9). Thus, plasma sprayed deposits consist of individual splats, which are connected together by mechanical and chemical bonding. The inter-lamellae boundaries are associated with a significant level of micro cracks and pores (Ref 10). The physical and mechanical properties of the vacuum-plasma-sprayed structures are lower than those of conventionally processed materials as the result of their internal microstructure. This is due to the existence of about 3% porosity, interlamellar boundaries, and internal defects. Therefore, to achieve the desired level of mechanical performance from the VPSF Ti-6Al-4V components, the level of porosity and internal defects must be reduced, interlamellar boundaries must be eliminated, and optimum phase compositions must be achieved through post deposition heat treatments (Ref 7).
Since microstructure produced by this process is completely different from microstructure produced by casting and forging processes, it is expected to have deferent response to heat treatment. To achieve one or more appropriate heat treatment processes for such microstructures, it is needed to study the kinetic behavior of this structure at high temperature. In this paper, the kinetic of transformations in VPSF Ti-6Al-4V alloy (α+β Ti alloy) was studied, comprehensively.
Materials and Methods
Ti-6Al-4V alloys are classified as α+β alloys that martensitic transformations in these alloys occur during quenching. In the present study, βα+β transformation kinetics on the sprayed and homogenized sample at 1000oC in β phase was evaluated. The spray-formed Ti-6Al-4V samples were made by deposition of standard Ti-6Al-4V feedstock powders using a vacuum plasma spraying process. There are important parameters effecting The VPSF process such as plasma gas composition and flow rate, pressure, current, Spray distance and etc. More details about theses parameters were described in H.R. Salimijazi et al. work (Ref 7). The as-sprayed samples were homogenized in β phase followed by fast cooling to two phase temperature region (800, 850, 900 and 950oC). Finally, after formation of α+βphases, samples were quenched in water to suppress further transformation. Heat treatment tests were performed in a tube furnace under control argon atmosphere.
The heat-treated specimens were cross-sectioned, epoxy mounted, polished, and etched in Kroll’s solution. To observe the microstructure of the polished and etched specimens, a Nikon microscope equipped with Sony DSC-H50 camera was used for optical microscopy measurement of the equilibrium phases and image analysis was performed using Clemex software.
Results and Discussions
Heat Treatment
After heat treatments of the VPSF Ti-6Al-4V alloy, and βα+β phase transformation, the amount of α phase obtained at different isothermal temperatures and times was measured by image analysis.
The level of α phase was increased by increasing the transformation time (Fig 1). Decreasing the transformation temperature results in increasing the degree of undercooling which plays a key role as transformation stimulant. The more the transformation stimulant, the faster the transformation. Figure 1: Kinetics of the βα+β transformation at different
temperatures in isotherm state for VPSF Ti-6Al-4V alloy. SEM images before and after heat treatments showed that the structure completely changed from splat shape to grain structure and morphology of pores converted to spherical shaped porosities.
Figure 2:SEM images from VPSF Ti-6Al-4V alloy before (a) and after (b) the heat treatment
Metallography
Microstructures of the VPSF Ti-6Al-4V samples after the isothermal heat treatments at various temperatures and times are shown in Fig. 3and 4. At a constant transformation temperature, by increasing the transformation time, the amount of α phase increase d from 27.8% to 34.8%. By
20μm
gradually increasing the transformation time, grain boundaries were covered with nuclei.
Figure 3: Microstructure of the heat-treated VPSF Ti-6Al-4V alloy at 800oC for (a) 10s, (b) 20s, (c) 40s, (d) 60sec.
Microstructure of the Ti-6Al-4V alloy at transformation temperatures from 800 to 950oC and constant transformation time of 60 sec is shown in Fig. 4. The number of nucleation sites was higher in heat treatment temperatures of 800 and 900 oC, as seen in Fig. 4 (a) and (b). T he nuclei of α phase were located in BOTH grain boundaries and inside the grains. It seems that these nucleation sites were distributed within the β grain homogeneously. At transformation temperature above 900oC, nucleation sites were limited to the grain boundaries as seen in Fig. 4(d) and (c).
Figure 4: Microstructure of heat-treated VPSF Ti-6Al-4V alloy at different temperature of (a) 800oC, (b) 850oC, (c) 900oC, (d) 950oC.
As the transformation temperature increased from 800oC to 950oC, the percent of α phase decrease d from 34.8% to 29.7%.
Transformation Kinetic
Phase transformations are of central importance in materials science and engineering. An understanding of the thermodynamics of phase equilibrium is the foundation for understanding their kinetics (Ref 11). For many solid-state transformations, the rate of transformation is sufficiently slow that it can be followed in time. The progress of these transformations can be described by the Johnson –Mehl –Avrami equations (Ref 12). For the isothermal transformations this equation has the following popular form (Ref 2, 3):
where, y represents the product phase volume fraction which varies with time t sec, k the reaction rate constant and n Avrami index that describes the nucleation and growth mechanisms (2,3). In the present study the kinetics of the β α+β phase transformat ion in VPSF Ti-6Al-4V alloy were modeled by adapting the classic JMA Johnson-Mehl-Avrami theory.
It must be pointed out that the JMA equation traces the transformation from its start and does not describe the pre-processing stages and the incubation time. Hence, the time should be treated not as an absolute time but as relative to the start of the transformation.
Eq. (1) was used to analyze the experimental data by means of
logarithmic plots, where
is plotted versus ln(t).
The slope of the resulting straight line is the Avrami exponent n, while from the intercept the k value can be calculated. Such plot is presented in Fig. 5 for the temperatures applied in the present study. The n and k values for all applied temperatures can be derived for the alloy from Fig. 5. The extracted data were tabulated in Table 1.
Figure 5 : Plots of
against ln(t) for deriving of
the JMA parameters for VPSF Ti –6Al –4V alloy at different temperatures.
a b
c d
60μm
a b
c d
100μm
100μm 100μm
100μm
It is clear that for lower temperature (800 and 850oC), the experimental measurements was well described by a single straight line. This theory can be used to describe the kinetics of the phase transformation βα+β in the Ti-6Al-4V alloy at high temperatures. In addition, the transformation mechanism doesn’t change during the transformation path.
At higher temperatures (900 and 950oC) there is a tendency to change line slope. The data consisted of two parts that can be divided into two separate lines. The explanation of these observations is that the transformation mechanism during the transformation path has been changed. The amount of k or transformation rate constant of the VPSF Ti-6Al-4V alloy was much larger than the alloy produced by casting technique. This can be contributed to the higher rate of nucleation in the plasma sprayed microstructure. As the result of defects, ultra fine grain and high level of residual stresses in VPSF Ti-6Al-4V component compare to those in casting specimen. As the result, they can increase the stimulant of transformation and therefore achieved a higher transformation rate.
Table 1: JMA kinetic parameters for Ti–6Al–4V alloy obtained by versus ln(t) plots.
Mechanism of the phase transformation at temperatures lower than 900oC, was including homogeneous nucleation and growth of the α-phase. Lower transformation temperatures are equivalent to higher degree of undercooling. Thus, the transformation stimulant in these temperatures was higher. In this condition, nucleation essentially homogeneous rate controls the transformation rate. But the transformation mechanism at temperatures above 900oC was variables and included in two parts. The mechanism of first part was formation of the α-phase in the grain boundary. And mechanism of the second part is α-plate nucleation and growth from the grain boundaries. By increasing the temperature, the transformation rate constant was also increased. This shows that the overall rate of transformation in this temperature range is controlled by diffusion.
Nucleation and Growth
At lower temperatures, due to the slower diffusion and larger undercooling ↑ΔT=,nucleation stimulant was higher and the nucleation rate would be higher than the growth rate. Therefore, the formed α- phase from the βα+β phase transformation, due to the slower diffusion, had lower growth and the final microstructure was finer.
At higher temperatures, due to the faster diffusion and smaller undercooling ↓ΔT=, nucleation stimulant was lower and growth rate would be higher than the nucleation rate. Therefore, the formed α-phase from the βα+β phase transformation has a higher growth and the obtained microstructure was coarser.
Figure 6: Microstructure of the heat-treated VPSF Ti-6Al-4V alloy at deferent temperature and time: (a,b,c,d) at 800oC (<900oC) and 10, 20, 40, 60 sec and (e,f,g,h) at 950oC (>900oC) and 10, 20, 40, 60 s.
The number of nuclei at lower temperatures ranges from 800 to 850oC was more than higher temperatures ranges from 900 to 950oC, as shown in Fig. 6. At high temperatures, the transformation rate controller is nucleation, which has the lowest rate. It is assumed that the nucleation rate directly depends on the degree of undercooling, . Conversely, at low temperatures, the deposition kinetics is controlled primarily by the growth rate that changes with the absolute temperature.
Time-Temperature-Transformation Diagram
An isothermal transformation diagram for the kinetics of βα+β phase transformation in VPSF Ti-6Al-4V alloy was a
d
g
c
f
b
e
h
60μm
60μm
60μm
60μm60μm
60μm60μm
developed by using the kinetic results at different constant temperatures Fig. 7. The start and end of the transformation is difficult to be determined reliably. It should be mentioned that the end of transformation does not correspond to the situation where the initial phase (β-phase) completely transforms to the new phase (α-phase). At each temperature, a different phase composition α+β mixture is the final product. For this alloy no nose point for the start of the transformation was found. Since the incubation periods for the studied transformations were very short (a few seconds) this effect may contribute to the experimental absence of the nose point. The developed TTT diagram can be used in the heat treatment practice for VPSF Ti-6Al-4V alloy.
Figure 7:Calculated time–temperature–transformation diagrams for VPSF Ti-6Al-4V alloy.
Summary and Conclusion
The βα+β transformation kinetic in VPSF Ti–6Al–4V alloy was studied at isothermal conditions using metallography and image analysis. Different mechanisms of the transformation were suggested. Mechanism of phase transformation at temperatures lower than 900oC included homogeneous nucleation and growth of α-phase. On the other hand, the transformation mechanism at temperatures above 900 oC was variable and occurred in two parts; the formation of α-phase grain boundary a nd α-plate nucleation and growth from the grain boundaries. The kinetics of the βα+βtransformation was modeled under isothermal conditions in the theoretical frame of the Johnson–Mehl–Avrami (JMA) theory. The Avrami index, n, and the transformation rate constant, k, were determined. Finally, the Time–temperature–transformation diagram was developed for the βα+βtransformation for VPSF Ti–6Al–4V alloy. Iso-lines were calculated and plotted for tracing the amounts of the α-phase.
References
1. S. Tamirisakandala, B.V. Vedam, R.B. Bhat, Recent
Advances in the Deformation Processing of Titanium
Alloys, Journal of Materials Engineering and
Performance, 2003, Vol. 12, p 661-673.
2. S. Malinove, P. Markovsky, W. Sha, Resistivity Study
and Computer Modeling of the Isothermal Transformation
Kinetics of Ti-8Al-1Mo-1V Alloy, Journal of Alloys and
Compounds, 2002, Vol. 333, p 122-132.
3. S. Malinove, P. Markovsky, W. Sha, Z. Guo, Resistivity
Study and Computer Modeling of the Isothermal
Transformation Kinetics of Ti-6Al-4V, Ti-6Al-2Sn-4Zr-
2Mo-0.08Si Alloys, Journal of Alloys and Compounds,
2001, Vol. 314, p 181-192.
4. L.M. Gammon, R.D. Briggs, J.M. Packard, K.W. Batson,
R. Boyer, C.W. Domby, Metallography and
Microstructures of Titanium and Its Alloys, Metallography
and Microstructures, , ASM Handbook, ASM International,
2004, Vol. 9, p 899–917.
5. D.M. Stefanescu, R. Ruxanda, Solidification Structures of
Titanium Alloys, Metallography and Microstructures,
ASM Handbook, ASM International, 2004, Vol. 9, p 116–
126.
6. H.R. Salimijazi, T.W. Coyle, J. Mostaghimi, Vacuum
Plasma Spraying: A New Concept for Manufacturing Ti-
6Al-4V Structure, JOM, 2006.
7. H.R. Salimijazi, T.W. Coyle, J. Mostaghimi,Understanding
Grain Growth and Pore Elimination in Vacuum-Plasma-
Sprayed Titanium Alloy, Metallurgical and Materials
Transactions A, 2007, Vol. 38A, p 476-484.
8. K. Balani, A. Agarwal, T. Mckechnie, Near Net Shape
Fabrication via Vacuum Plasma Spray Forming,
International Symposium Research Students on Material
Science and Engineering, 2004.
9. Fauchais P., Understanding Plasma Spraying, J. Phys. D:
Appl. Phys., 2004, Vol. 37, p 86-108.
10. H.R. Salimijazi, M. Raessi, J. Mostaghimi, T.W. Coyle,
Study of Solidification Behavior and Splat Morphology of
Vacuum Plasma Sprayed Ti Alloy by Computational
Modeling and Experimental Results, Surface and Coatings
Technology 201, 2007, p 7924-7931.
11. W. Robert Balluffi, M. Samuel Allen, W. Craig Carter,
Kinetics of Materials, John Wiley & Sons, Inc., Hoboken,
New Jersey, 2005.
12. A. Kenneth Jackson, Kinetic Processes, Wiley-VCH
Verlag GmbH & Co. KGaA, Weinheim, 2004
ββ
Structure and Properties of HVOF and Plasma Sprayed Ceramic Alumina-Chromia Coatings Deposited from Fused and Crushed Powders P. Vuoristo*, K. Niemi, V. Matikainen, L. Hyv?rinen, H. Koivuluoto
Tampere University of Technology, Department of Materials Science, Tampere, Finland
*E-mail: petri.vuoristo@tut.fi
L.-M. Berger
Fraunhofer Institute for Material and Beam Technology (IWS), Dresden, Germany
S. Scheitz, I. Shakhverdova
Dresden University of Technology, Dresden, Germany
Abstract
Different feedstock powder compositions of the alumina-chromia system (75/25, 60/40 and 25/75 in wt.-%) were deposited by conventional plasma spraying, 3-anode plasma spraying and by HVOF spraying. The powders for the plasma spray processes had particle sizes of -38+10 ?m and for HVOF spraying -25+5 ?m and -25+10 μm. The coatings were evaluated by their microstructure, phase composition, wear and corrosion properties as well as electrical properties. The study showed that the effect of the spray process is crucial for the wear properties and revealed that particularly the HVOF spray process results in coatings with dense structures and excellent wear properties. Despite of still high porosity 3-anode plasma coatings show higher hardness than conventional APS coatings and can be sprayed with higher feed rates. The coating properties appear not to have a linear dependence on the chromia content.
Introduction
Both alumina (Al2O3) and chromia (Cr2O3) are widely used for preparation of thermally sprayed coatings. These coatings are predominantly deposited by atmospheric plasma spraying (APS). High velocity spray methods, such as high velocity oxy-fuel (HVOF), are still seldom used due to lower feeding rates and deposition efficiency. However, coatings deposited by HVOF have been reported to give greatly improved wear resistance (Ref 1). These coatings can be utilized e.g. in applications where the spraying time is relatively short or coating thickness is low. Also these coatings can fulfill demands that are normally obtained only by more expensive materials and manufacturing processes.
Corundum (?-alumina) and the isostructural eskolaite (Cr2O3) form solid solutions that can be designated as (Al,Cr)2O3. At high temperatures, the formation of a solid solution is possible for the full range of concentrations; its decomposition due to an existing miscibility gap in the phase diagram depends on the reaction kinetics.
Both Al2O3 and Cr2O3 show also a specific material behavior in the thermal spray processes (Ref 2). The transformation of the thermodynamically stable corundum (?-alumina) and the formation of metastable phases, such as ?-alumina, is widely discussed in the literature. Stabilization of ?-alumina by alloying with chromia has already widely been studied (Ref 3-10). A stabilization can be obtained by using mechanically blended Al2O3 and Cr2O3 feedstocks and special spray processes, such as water stabilized plasma (WSP) (Ref 5). However, the use of (Al,Cr)2O3 solid solution feedstock particles is more effective approach (Ref 5).
Main problem of spraying chromia is the high volatility and associated low deposition rate, which is in connection with the high melting point particularly critical for HVOF (Ref 2). For this reason powder compositions Cr2O3-25 wt.-% TiO2 are used, but titania alloying lowers the hardness and wear resistance of the coatings too much. In order to improve the deposition efficiency and to maintain the coating properties compared with pure chromia, alloying with alumina is regarded as a better choice (Ref 11).
Thus, the disadvantages of Al2O3 and Cr2O3 can be at least minimized by addition of the other oxide. However, hardness and abrasion wear resistance of APS-sprayed coatings show no linear behavior in dependence of the chromia content (Ref 12).
In this work a range of compositions (75/25, 60/40 and 25/75) of the Al2O3-Cr2O3-systems has been studied with particular focus on the sprayability of the feedstock powders by HVOF
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and 3-anode plasma. Comparison between the HVOF and
plasma processes was carried out.
Experimental
Powders
Fused and crushed powders used in the present study were
supplied by Ceram GmbH (Albbruck-Birndorf, Germany) and are listed in Table 1. Three different compositions Al 2O 3/ Cr 2O 3 75/ 25, 60/ 40, and 25/ 75 (in wt.-%) were studied. Different particle sizes were available; these were -38+10 μm, -25+5 μm and -25+10 μm. The coarsest powder was used in the plasma spray and the two finer powders in the HVOF spray process.
Table 1: Powders used for APS and HVOF spraying at TUT (1)
and IWS (2) and particle size, values of d 90, d 50 and d 10 as given by the producer.
Al 2O 3/ Cr 2O 3 Particle size d 90 [μm] d 50 [μm] d 10
[μm]
25/ 75 -38+10 μm (1), (2) 45.3 25.8 10.9
-25+5 μm (1), (2) 23.6 15.8 7.1
-25+10 μm (1)
27.0 16.8 9.1
60/ 40 -38+10 μm (1), (2)
41.5 24.8 12.4
-25+5 μm (1), (2)
22.4 15.6 7.6
-25+10 μm (1)
25.6 16.9 10.2
75/ 25 -38+10 μm (1), (2)
43.9 24.6 12.4 -25+5 μm (1), (2)
23.1 14.5 7.3 -25+10 μm (1) 27.5 17.5 10.0
Spray Processes
The powders were sprayed by APS at Tampere University of
Technology (TUT) using a Sulzer Metco A3000 system with
F4 plasma gun (Ar/H 2 plasma) and at Fraunhofer IWS by
using a F6 Plasma gun and a 3-anode Delta gun (both Ar/H 2
plasma; GTV mbH, Luckenbach, Germany). In both
organizations, HVOF spraying was performed using a
TopGun (GTV mbH) with ethene (C 2H 4) as fuel gas. The
spray parameters are listed in Table 2.
Powder and Coating Characterization
The morphology of the feedstock powders was analyzed by
SEM (XL30, Philips). The measurement of particle size was
made with laser diffraction (CILAS 990LD, Quantachrome
GmbH & Co. KG).
The microstructure was analyzed by conventional optical
microscopy using metallographically prepared cross sections.
The phase composition of the coatings as well as the powders
was determined by X-ray diffraction (XRD) using a Bruker
AXS D8 Advance diffractometer with a CuK D radiation.
Vickers hardness (HV 0.3) was measured from the polished
cross sections of the coatings (MMT-X7, Matsuzawa at TUT
and HP-mikromat 1-HMV, Hegewald & Peschke Me?- und
Prüftechnik GmbH at IWS). The results are given as an
average of ten measurements. The abrasion wear resistance was studied at TUT with a device modified from ASTM G65. Dry quartz sand (SiO 2) with a grain size of 0.1-0.6 mm was used as an abrasive. Table 2: Parameters for APS and HVOF processes at TUT (1) and IWS (2). Process Parameter Al 2O 3/ x Cr 2O 3 x = 25 x = 40x = 75
H V O F
TopGun (1) C 2H 4[slpm] 93 93 93
O 2 [slpm] 272 272 272 Nozzle [mm] 135 135 135 TopGun (2)
C 2H 4[slpm] 90 90 90
O 2 [slpm] 245 245 245
Nozzle [mm] 135 135 200
A P S F4(1)
Ar [slpm] 41 41 38 H 2[slpm] 14 14 13 Power [kW] 43 44 44 F6(2)
Ar [slpm] 40 40 40 H 2[slpm] 13 8 8 Power [kW] 62 51 51 Delta (2) Ar [slpm] 40 40 40 H 2[slpm] 8 8 8 Power [kW] 67 62 67
Corrosion exposure tests were carried out for evaluation of the
long-term chemical stability of the coatings in acidic (pH 2) and basic (pH 13.2) solutions. The coatings were removed from the steel substrates as fragments and cleaned in acetone using ultrasonic bath. The detached fragments were dried in an oven, cooled in a desiccator and weighed. The fragments were placed in plastic vessels containing 200 ml of the exposure solution with desired acidity. The pH 2 exposure liquid was prepared to ion exchanged water using H 2SO 4 and the pH 13.2 exposure liquid to ion exchanged water using NaOH. After exposure tests (310 h at 85 °C), the solution was analyzed using ICP-AES (Inductively Coupled Plasma-Atomic Emission Spectrophotometry) by Eurofins Scientific Finland Oy and the contents of dissolved aluminum and chromium were determined. The electrical resistivity was analyzed at IWS using an impedance spectrometer IM6ex (Zahnerelektrik GmbH & Co. KG, Germany) and the same procedure as described previously (Ref 13). The surface of the coatings was grinded before the measurement. Results and Discussion Powder Characterization Figure 1 shows an example of the typical irregular shape of fused and crushed powders as studied by SEM. The morphologies were the same for all compositions. Powders
with a particle size of -25+10 μm was found to have somewhat better sprayability in the HVOF process compared to powders with a particle size of -25+5 μm. The finer fraction caused slightly unstable powder feeding during spraying, due to presence of fine particles in these powders. The particle sizes (d 90, d 50 and d 10) are compiled in Table 1 above.
Figure 1: SEM (BSE) image of an Al 2O 3/ 75 Cr 2O 3
powder with a particle size of -25+10
μm.
Particles appearing white in the micrograph are metallic
chromium according to the EDS analysis, which is detectable also in the XRD pattern (later). Also, variations in the grayscale are due to differences in the composition of the
particles consisting of different amounts of Cr 2
O 3
and Al 2
O 3
.
Coatings Microstructures
Figure 2 shows the optical micrographs of all coatings
prepared in this study. The APS-sprayed coatings (F4, F6 and Delta) have a different porosity. The highest porosity was
observed in Delta sprayed coatings. The HVOF-sprayed TopGun coatings, especially those prepared at TUT, showed
the highest density. The coating microstructures for each spray process were nearly independent from the Cr
2O
3 content. It
appears also that the amount of bright stripes is higher in the APS coatings. Coating thicknesses in Fig 2 are not to be compared between processes as the powder feeding and
deposition parameters were not exactly identical at IWS and TUT.
Figure 3 shows as an example the diffraction pattern of the 75/25 coating sprayed with the Delta process. The coatings consist mainly of ?- and ?-phases, which composition is a solid solution of chromia and alumina (Al,Cr)2O 3. Besides these two phases, metallic chromium was observed by XRD and it is present as white lamellae in the cross sections (Fig. 2). Presence of the metallic chromium was highest for the 75/25 and 60/40 compositions. In APS coatings (F4, F6 and Delta) the white zones were larger and the Cr content was found up to 6 wt.-%, while less than 2 wt.-% Cr content was found in HVOF coatings.
25 wt.-% Cr 2O 3 40 wt.-% Cr 2O 3 75 wt.-% Cr 2O 3
T o p G u n (1)
T o p G u n (2)
F 4 (1)
F 6 (2)
D e l t a (2)
Figure 2: Coating microstructure for the different spray processes at TUT (1) and IWS (2), (the length scale is valid for all micrographs).
Figure 3: XRD of a Delta sprayed Al 2O 3/ 25 Cr 2O 3 coating.
Table 3 shows a comparison of the used powders and sprayed coatings and the content of Cr 2O 3 in solid solution (Al,Cr)2O 3 as determined by XRD. Depending on the spray processes, the content of Cr 2O 3 in solid solution is different between the coatings as well as the used powders.
10 μm
2Z [°] CuK D
wP
Table 3: Comparison of Cr 2O 3 content in solid solution for powders and coatings sprayed by TUT (1) and IWS (2) from XRD analysis ( can not be determined, too many side peaks).
Cr 2O 3 content /
wt.-%
Al 2O 3/ x Cr 2O 3
x = 25 x = 40 x = 75 Powders 20 (1) (1) (1) 20 (2) 40 (2) 79 (2)
TopGun (1)
- - 70 TopGun (2)
20 39 76 F4 (1) 16 39 59 F6 (2) 23 36 76 Delta (2)
21 37 76
Microhardness of the Coatings
Figure 4 shows a comparison of the coating hardness. With increasing amount of Cr 2O 3, the hardness increases. The highest values were reached at HVOF-sprayed 25/75 composition (1400 HV 0.3). Difference in hardness between the HVOF spayed coatings from TUT and IWS may result from the different particle size (-25+10 μm at TUT and -25+5 μm at IWS). Despite of the high porosity of the Delta sprayed coatings, they showed a high hardness, somewhat similar to the HVOF sprayed coating (e.g. 60/40 composition). However, Delta sprayed coatings did not reach the hardness of TopGun sprayed 25/75 coatings but the deposition rate (thickness per pass with similar surface speed) was nearly ten times higher. This coating has also a high hardness of 1200 HV 0.3. F4 and F6 sprayed coatings showed lower hardness values.
Figure 4: Microhardness (HV 0.3) of APS (F4, F6, Delta) and HVOF (TopGun) sprayed Al 2O 3/ Cr 2O 3 coatings at TUT (1) and IWS (2).
Wear Resistance of the Coatings
Figure 5 shows a comparison of the mass loss in the abrasion wear test for the coatings sprayed at TUT. In the case of the HVOF coatings, the mass loss is practically independent from the Cr 2O 3 content. The mass loss of the APS sprayed coatings is significantly higher. The ratio of weight loss between the APS and HVOF sprayed Al 2O 3/Cr 2O 3 coatings was about
nearly 12:1, 6:1 and 11:1 for 75/25, 60/40 and 25/75
compositions..
HVOF (TopGun) Al 2O 3/ Cr 2O 3 coatings sprayed at TUT.
Corrosion Properties
Figures 6 and 7 show the amounts of dissolved aluminum and chromium in the coatings. Improved corrosion stability in the acidic (pH 2) and basic (pH 13.2) test conditions was obtained for Al 2O 3/Cr 2O 3 coatings when the dissolution of aluminum was significantly reduced. The stability was improved as a function of the chromia content as shown in Fig. 7. Exposure tests showed that also dissolution of chromium was low with these coatings as shown in Fig. 8.
Figure 6: Dissolution of aluminum from APS (left) and HVOF (right) Al 2O 3/Cr 2O 3 coatings at two pH conditions (pH 2 and pH 13.2) at 85 °C for 310 h.
H a r d n e s s H V 0.3
Cr 2O 3
12
5.7 1.5
11
4.5
1.6
29
14
9.3
24
7.9
2.9
510
152025303575/25
60/40
25/75
75/25
60/4025/75
A l [m g /l ]
Al 2O 3/Cr 2O 3[wt.-%]
pH 2pH 13.2
TopGun (HVOF)
F4 (APS)
Figure 7: Dissolution of chromium from APS (left) and HVOF (right) Al 2O 3/ Cr 2O 3 coatings at two pH conditions (pH 2 and pH 13.2) at 85 °C for 310 h.
Alumina coatings show a good corrosion resistance in the pH range of 4 to 10. Outside this range, the corrosion resistance decreases and dissolution of alumina controls the overall corrosion of the coatings (Ref 1). Earlier similar corrosion exposure tests have shown that dissolution of aluminum from plain Al 2O 3 coatings deposited by APS and HVOF was significantly higher, of the order of 130 mg/l at pH 2 and 290 mg/l at pH 13.2 for the APS coatings (Ref 8). As chromia alloying improves the corrosion resistance of alumina coatings it gives broader pH range where these coatings can be utilized.
Electrical Properties
Al 2O 3 coatings are widely used for electrical insulating applications. But the content of Cr 2O 3 will be influencing the resistivity of the coatings negatively. The spectra in Fig. 8 shows the resistivity of the coatings sprayed by APS (F6). The values of resistivity were calculated from the measured impedance spectrum (Ref 13). The resistivity of the coatings decreased from 2.1x1010 Ohm m to 2.1x108 Ohm m. Thermally sprayed coatings of pure Cr 2O 3 have a resistivity of nearly 106 Ohm m (Ref 14) and of pure Al 2O 3 of nearly 1011 Ohm m (Ref 12, 15). There is no linear dependence of the resistivity on the Cr 2O 3 content. The 60/40 coatings had a higher resistivity than the 25/75 and 75/25 coatings. The effects of the coating microstructure and the amount of metallic chromium on the resistivity have to be clarified in further experiments.
Figure 8: Impedance spectrum of F6-sprayed Al 2O 3/Cr 2O 3 coatings, calculated values of resistivity based on an R//C model are given in square brackets in the legend.
Summary and Outlook
Both plasma spraying and HVOF spraying were used in this study to prepare coatings of three compositions of the Al 2O 3/Cr 2O 3 system. The results can be summarized as follows:
x All coatings sprayed by HVOF and APS processes from
powders of the Al 2O 3/Cr 2O 3 system were found to have good microstructures and properties; x
The HVOF sprayed coatings showed denser structures and much improved wear properties compared with the APS sprayed coatings;
x
The HVOF sprayed coatings were significantly higher in their hardness, particularly due to better coatings microstructures;
x
Removal of fine particles from -25+5 μm size powders improved significantly their processability by HVOF; for this reason a particle size of -25+10 μm was used for HVOF experiments at TUT resulting in coatings with highest hardness and excellent abrasion wear resistance; x
APS sprayed coatings in general had slightly higher porosity, but it was found that the Delta APS sprayed coating reached hardness values close to those of the HVOF coatings, but with a higher deposition efficiency than with HVOF processes;
x
There are bigger influences of resistivity than the content of Cr 2O 3; i.e. the amount of chromium could decrease the value of resistivity.
In further experiments the corrosion and wear resistance of Delta sprayed coatings will be investigated. Also the electrical properties of the Delta and TopGun sprayed coatings will be studied in more detail.
10
10
10
10
10
10
10
10
10
10
10
10
101010101010101010101010R e s i s t i v i t y / O h m m
Frequency / Hz
3.1
2.6
1.4
3.9 2.8 1.1
5.4
5.8
3.8
3.5
2.5
2.5
1234
56775/25
60/40
25/75
75/25
60/40
25/75
C r [m g /l ]
Al 2O 3/Cr 2O 3[wt.-%]
pH 2
pH 13.2
TopGun (HVOF)
F4 (APS)
Acknowledgement
Results of Dresden University of Technology presented here are part of the research project ‘New ceramic coating powders for thermal sprayed coatings, (‘Neue keramische Beschichtungspulver für thermisch gespritzte Schichten) contract KF 2097505KI9, which is funded via AiF by the Federal Ministry of Economics and Technology. Results of Tampere University of Technology were partly funded by the High-ALPHA project, which was funded by Tekes – the Finnish Funding Agency for Technology and Innovation, Tampere University of Technology, and by industry. The authors gratefully acknowledge these funding organisations for their support. The authors would also like to thank B. Wolf (IWS) for preparation of the cross sections and hardness measurements, and Mr. Mikko Kylm?lahti, of Tampere University of Technology, Department of Materials Science, for the spraying of the coatings at TUT.
References
1. K. Niemi K, Abrasion Wear Characteristics of Thermally
Sprayed Alumina Based Coatings. Doctoral Thesis, Tampere University of Technology, Publication 820, Tampere 2009, ISBN 978-952-15-2186-7, p. 117.
2. L.-M. Berger, C.C. Stahr, F.-L. Toma, G.C. Stehr, and E.
Beyer, Ausgew?hlte Entwicklungstendenzen bei der Herstellung thermisch gespritzter keramischer Schichten, Jahrbuch Oberfl?chentechnik 2007, Vol. 63, Ed.: R.
Suchentrunk, Eugen G.Leuze Verlag, Bad Saulgau, 2007, p. 71-84 (in German)
3. P. Chráska, J. Dubsky, K. Neufuss, and J. Písacka,
Alumina-Base Plasma-Sprayed Materials Part I: Phase Stability of Alumina and Alumina-Chromia, J. Thermal Spray Technol., 6 (1997), Issue (3), p 320-326
4. B. R. Marple, J. Voyer, and P. Béchard, Sol Infiltration
and Heat Treatment of Alumina Chromia Plasma-Sprayed Coatings, J. Eur. Cer. Soc., 2001, 21 (7), p 861-868
5. C.C. Stahr, S. Saaro, L.-M. Berger, J. Dubsky, K. Neufuss,
K. and M. Herrmann, Dependence of the Stabilization of D-Alumina on the Spray Process, J. Therm. Spray Technol.
16 (2007), Issue (5-6), pp. 137-147.
6. J. Dubsky, P. Chraska, B. Kolman, C.C. Stahr, and L.-M.
Berger, Phase Formation Control in Plasma Sprayed Alumina-Chromia Coatings. Ceramics - Silikaty55 (2011), Nr.3, p.294-300
7. K. Yang, X. Zhou, H. Zhao, and S. Tao, Microstructure
and Mechanical Properties of Al2O3-Cr2O3 Composite Coatings Produced by Atmospheric Plasma Spraying., Surf. Coatings Technol., 2011, 206(6), p 1362-1371
8. K. Niemi, J. Hakalahti, L. Hyv?rinen, J. Laurila, P.
Vuoristo, L.-M. Berger, F.-L. Toma, and I. Shakhverdova, Influence of Chromia Alloying on the Characteristics of APS and HVOF Sprayed Alumina Coatings, Proc. Int.
Thermal Spray Conf, Hamburg, Sept. 27 - 29, 2011: DVS Media, (DVS-Berichte 276), 2011, p 1179-1184 9. L. Hyv?rinen, J. Hakalahti, M. Kylm?lahti, J. Silvonen, K.
Niemi and Vuoristo, P. Improving the Properties of
Plasma and HVOF Sprayed Alumina Coatings by Chromia
Addition. Thermal Spray 2012: Proc. Int. Thermal Spray
Conf. May 21-24, 2012, Houston, Texas, USA: ASM
International, 2012, p. 488-493
10. K. Yang, J. Feng, X. Zhou, and S. Tao, Microstructural
Characterization and Strengthening-Toughening Mechanism of Plasma-Sprayed Al2O3-Cr2O3 Composite
Coatings, J. Therm. Spray Technol., 2012, 21 (5), p. 1011-
1024
11. S.H., Yu, H. Wallar, Chromia spray powders, Saint-
Gobain Ceramics and Plastics, Inc., 2006, U.S. Patent No.
7012037
12. T. B?rner, L.-M. Berger, S. Saaro, and S. Thiele,
Systematische Untersuchung der Abrasionsbest?ndigkeit
thermisch gespritzter Schichten aus dem Werkstoffsystem
Al2O3-TiO2-Cr2O3, Tagungsband zum 13.
Werkstofftechnischen Kolloquium in Chemnitz, 30.
September - 01. Oktober 2010, Schriftenreihe "Werkstoffe
und werkstofftechnische Anwendungen", Band 37,
Herausgeber: B. Wielage. Chemnitz: TU Chemnitz, 2010,
p. 178-187. (in German)
13. F.-L. Toma, L.-M. Berger, S. Scheitz, S. Langner, C.
R?del, A. Potthoff, V. Sauchuk and M. Kusnezoff,
Comparison of the Microstructural Characteristics and
Electrical Properties of Thermally Sprayed Al2O3 Coatings
from Aqueous Suspensions and Feedstock Powders, J.
Thermal Spray Technol,21 (2012), Issue (3-4), p. 480-
488.
14. L.-M. Berger, S. Saaro, C.C. Stahr, S. Thiele and M.
Woydt, Entwicklung keramischer Schichten im System
Cr2O3-TiO2., Thermal Spray Bull.,2009, 2(1), p 64-77
15. F.-L. Toma, S. Scheitz, L.-M. Berger, V. Sauchuk, M.
Kusnezoff, and S. Thiele, Comparative Study of the
Electrical Properties and Microstructures of Thermally
Sprayed Alumina- and Spinel-Coatings, J. Thermal Spray
Technol., 20 (2011), Vol., Issue 1-2, p. 195-204.
Structural Modifications on Material Surfaces by Thermal Nanoparticle Spraying and Microlines Patterning
S. Kirihara
JWRI, Osaka University, Ibaraki, Osaka, Japan
*E-mail: kirihara@jwri.osaka-u.ac.jp
Abstract
Thermal nanoparticles coating and microlines patterning were newly developed as novel technologies to fabricate fine ceramics layers and geometrical intermetallics patterns for mechanical properties modulations of practical alloys substrates. Nanometer sized alumina particles were dispersed into acrylic liquid resins, and the obtained slurries were sputtered by using compressed air jet. The slurry mists could blow into the arc plasma with argon gas spraying. On stainless steels substrates, the fine surface layers with high wear resistance were formed. In cross sectional microstructures of the coated layers micrometer sized cracks or pores were not observed. Subsequently, pure aluminum particles were dispersed into photo solidified acrylic resins, and the slurry was spread on the stainless steel substrates by using a mechanical knife blade. On the substrates, microline patterns with self-similar fractal structures were drawn and fixed by using scanning of an ultra violet laser beam. The patterned pure metal particles were heated by the argon arc plasma spray assisting, and the intermetallics or alloys phases with high hardness were created through reaction diffusions. Micro- structures in the coated layers and the patterned lines were observed by using a scanning electron microscopy.
Introduction
Thermal nanoparticles spraying had newly developed as a novel coating technique to create fine ceramic layers without structural defects on practical alloys substrates (Ref 1, 2). Thixotropic slurry of liquid resin including nanometer sized alumina particles was sputtered by using compressed air jet, and the slurry mist blew into the arc plasma with an argon gas spraying. In this investigation, fluid viscosity of the mixt pastes were measured and optimized to realize homogenized particles dispersions and smooth slurry flows. The relationship between the slurry properties and the formation of micro cracks and pores in the coated layers will be discussed. Subsequently, thermal microline patterning to realize metal phases drawing on the alloys substrates had been developed (Ref 3, 4). The liquid resin paste including micrometer sized aluminum particles were solidified as micro patterns by ultra violet laser scanning, and the intermetallic compound patterns were created through the heat treatment by the argon arc plasma spraying. In this investigation, microstructure and composite distributions in intermetallics patterns were observed and analyzed systematically. The reaction diffusion between the metal particles and substrate will be considered. Moreover, stress distributions on the metal surface were modulated by the geometric design of the intermetallics patterns. Load dispersion abilities will be discussed through numerical simulations and mechanical tests.
Experimental Procedure
The nanometer and micrometer sized alumina (TM-5D, Taimei Chemicals, Japan) and aluminum (ALE11PB, Kojundo Chemical, Japan) particles of 200 nm and 10 μm in average diameters were dispersed into the acrylic liquid resins at 40 %
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and 60 % in volume fractions, respectively. These obtained slurries were mixed in airtight containers through planetary configuration of rotation and revolution movements. Fluid characteristics of the slurry pasts were evaluated by using a viscosity and viscoelasticity measuring instrument to realize thixotropic fluid flows. The thixotropic slurry was injected into a gas flow chamber perpendicularly from a syringe nozzle of 40 mm in inner diameter, and the fluid viscosity was decreased by shear stress loading as schematically illustrated in Fig. 1. The slurry mists were sprayed toward a glass plate to measure and observe droplet diameters and spattering patterns. The micro sized mists were formed through the compressed air jet of 2 atm in gas pressure, and introduced into the into the arc plasma with an argon gas spray of 50 slpm in flow late through the stainless still guide nozzle of 4 mm in caliber size. Figure 2 shows the mist supplying equipment mounted on the plasma spraying apparatus. The alumina coated layer was formed on the SUS-304 stainless steel substrate of 50×50×1 mm in size placed at 140 mm in distance from the plasma gun. Crystal phase of the coated layer was analyzed by using X-ray diffraction method, and cross sectional microstructures and compositional distributions were observed by using an optical and scanning electron microscopy. Subsequently, self-similar patterns of Hilbert curve with stage numbers 1, 2, 3 and 4 of the fractal line structures were designed by using a computer graphic application, and these graphic images were converted into the numerical data sets by computer software. These patterns of 25×25 mm in whole size were composed of arranged lines of 400 μm in width. These graphic models were transferred into the processing apparatus as operating data sets. The pure aluminum particles were patterned on the SUS-304 stainless steel substrate by a stereolithographic pattering system as schematically illustrated in Fig. 3. The mixed resin paste with metal particles was spread with 100 μm in layer thickness on the substrate of 30×30×2 mm in size by a mechanically moved knife edge. An ultraviolet laser beam of
355 nm in wavelength and 100 μm in beam spot was
the resin surface. The solid pattern was obtained by light induced photo polymerization. Figure 4 shows the appearance of stereolithographic system. After the removing uncured resin by ultrasonic cleaning in ethanol solvent, the patterned sample was heated above the reaction temperature by using the argon gas arc plasma spraying. The microstructures and composite distributions were observed by using the scanning electron microscopy and energy dispersive X-ray spectroscopy. The stress distributions on the patterned samples were simulated by using the finite element method of the numerical simulation.
Results and Discussions
The acrylic liquid resin with the alumina nanoparticles of 200 nm in diameter at high volume contents exhibited the thixotropic characteristics of a non-Newtonian fluid. The formed slurry with the particles dispersion above 42 % in volume contents could not flow smoothly. For the stable materials supplying and the continuous mist creation, the maximum volume content of the nanoparticles can be optimized at 40 %. In this slurry fluid, the nanoparticles are considered to spatially disperse without coagglutinations. The obtained slurry blew toward the glass plate to observe and measure the mist droplets shapes and sizes. These droplet diameters became finer as shown in Fig. 5, and the variations could be reduced by decreasing the slurry supply rate at 130 cc/min. The minimum diameter size was 50 μm. Figure 6 shows the fine alumina coated layer formed on the stainless steel substrate. The ceramic layers of 50 μm in thickness were formed at 300 gpm in supply rate. The structural defects of cracks or pores were not observed through microscopic observations. The coated ceramics layer and alloy substrate were joined successfully without the defects of voids or exfoliations in the interfaces. Contamination of the carbon element in the created alumina layers derived from the acrylic liquid resins were not detected though the X-ray diffraction analysis. Moreover, dielectric constants of the coated layers had been measured as 8.6 by electromagnetic waves reflection spectroscopy. This value is lower than the ideal dielectric constant of 9.1. The alumina layer was considered to include the nanometer sized air cavities of 5 % in volume fraction approximately. Subsequently, the iron aluminide micro pattern with the fractal structure of Hilbert curve was formed successfully on the stainless steel substrate. Figure 7 shows the formed fractal polyline of the number 3 in fractal stage. The micrometer order geometric structure was composed of fine intermetallics lines of 450 μm in width. The part accuracy of these microline patterns were estimated as about 10 %. The iron aluminide composite was formed widely comparing with the designed line width though the metallic reaction diffusions.
Figure 8: A surface stress distribution on the Hilbert curve fractal pattern composed of the intermetallic compound visualized by using a numerical simulation of finite element method. The defects of crack or pores could not be observed in the cross sectional micro-structure of the formed intermetallics. The intermetallic layer and the alloy substrate were joined successfully without void formations through the reaction diffusion. The composite phase could be identified as Al-rich Fe3Al. During the heat treatment above 660 oC of the pure aluminum melting point, the reaction diffusion to synthesize the iron aluminide phase occurred between the pure aluminum particles and the steel substrate. After the solid phase reaction diffusion, the microstructure of the inter-metallics and alloys composites can exhibit about 1000 MPa in mechanical strength. The stress distributions on the patterned surfaces were visualized for the Hilbert curve of stage number 3 through the numerical simulation as shown in Fig. 8. The required mechanical properties of Young’s modulus were defined along the compositional analysis and the phase identifications. The stress intensities concentrate into the vicinity of fixed edge and are distributed along the patterned lines and the corner points with the higher hardness. The fractal patterns can include the more numbers of sides and nodes in limited aria comparing with the periodic arrangements of polygon figures.
Conclusion
Fine alumina coated layers and aluminide composite lines could be created successfully on stainless steel substrates by thermal nanoparticles spraying and microlines patterning to improve thermal and mechanical properties on components surfaces. As raw materials using these techniques, thixotropic slurries including nanometer and micrometer sized particles were formulate systematically. The slurry handling methods are key techniques to realize the nanoparticles arrangements create useful functional surfaces on various components. In the near future, the investigated coating and patterning techniques will become the candidate for efficient processes to realize effective thermal barrier coating without structural defects and mechanical properties improvement on various mechanical components.
References
1.S. Kirihara, Development of Thermal Nanoparticles
Spraying Technique, Thermal Spraying Technology,
2010, 30, p 44-50.
2.S. Kirihara, Introduction of Ceramics Nanoparticles into
Thermal Spraying, Journal of Japan Welding Society,
2010, 80, p 6-9.
3.Y. Uehara, S. Kirihara,Fabrication of Hard Alloys
Patterns with Fractal Structures on Light Metal Substrates through Reaction Diffusion, Journal of Smart
Processing, 2012, 1, p 186-189.
4.Y. Uehara, S. Kirihara,Modulation of Stress Gradients by
Intermetallics Patterning with Dendritic Fractals on Light
Metals, Journal of Functionally Graded Materials, 2012,
22
, p 36-42.
Electro-Catalytically Active Porous Nickel-Based Electrode Coatings Formed by Atmospheric and by Suspension Plasma Spraying
M. Aghasibeig, M. Mousavi, F. Ben Ettouill, R. Wuthrich, A. Dolatabadi*
Concordia University, Montreal, Quebec, Canada
*E-mail: ali.dolatabadi@concordia.ca
C. Moreau
National Research Council, Boucherville, Quebec, Canada
Abstract
Ni-based electrode coatings with enhanced surface areas, for hydrogen production, were developed using Atmospheric Plasma Spray (APS) and Suspension Plasma Spray (SPS) processes. The results revealed a larger electrochemical active surface area for the coatings produced by SPS compared to those produced by APS process. SEM micrographs showed that the surface microstructure of the sample with the largest surface area was composed of high amounts of small cauliflower-like aggregates with an average diameter of 10 μm.
Introduction
During the last century, increase of the global pollution and high costs of the fossil fuels has raised the demand of acquiring alternative resources of energy. Hydrogen as a green and renewable resource of energy is considered to be a promising alternative energy carrier to replace fossil derivative fuels. Alkaline water electrolysis is one of the most promising techniques for producing high purity hydrogen. An effort to obtain hydrogen by water electrolysis at diminishing costs with minimum energy consumption is to develop very active electrodes (Ref 1-3). Using highly active electrocatalysts with enhanced surface areas, for instance by forming porous electrode coatings, is a possible route of increasing the efficiency of the hydrogen evolution reaction (HER) (Ref 3, 4).
Although platinum is the best-known electrocatalyst for the HER, its use as cathode electrode material has been limited due to its expensiveness and rarity. In industrial applications, often nickel replaces platinum, since it is not only a catalytically active metal but also it has high stability in alkaline solutions at elevated temperatures (Ref 5).
Among different surface modification techniques, plasma spraying has demonstrated its capability to produce active porous electrode coatings, taking advantage of the surface roughness produced by the spraying process (Ref 4-6). The possibility of forming coatings at high deposition rates with moderate operating costs, and few limitations on the spraying materials and substrates has made this method superior compared to other surface coating techniques (Ref 7).
In recent years, there has been a major interest towards development of plasma sprayed coatings, utilizing nano-sized particles. This interest arouse from the unique features of nano-meter scale coatings, such as increased surface areas and superior performance. However, nano-particles cannot be used directly for plasma spraying due to the difficulties with their injection into the core of the plasma jet. Several methods have been developed to overcome the injection problem. One solution is using SPS, which is based on injection of a liquid feedstock. The suspension is basically made by dispersing the feedstock powder particles with the particle sizes ranging from tens of nanometers to a few micrometers in a liquid phase, which is mostly ethanol, water or a mixture of both. Addition of a dispersant is usually needed to reduce the agglomeration and sedimentation of the particles (Ref 8, 9).
The objective of this work was to develop porous Ni-based electrode coatings with enhanced surface areas for the hydrogen evolution reaction, using APS and SPS processes. The effects of spraying parameters of each process on the electrochemical active surface areas of the coated samples were studied. The results were compared for the coatings produced by APS and SPS processes.
Material and Methods
Spraying Materials
Two different commercial powders, nickel with the nominal particle size in the range of –75 to +45 μm (Metco 56C-NS) and nickel oxide with the nominal particle size in the range of of –5 μm to +500 nm (https://www.doczj.com/doc/7c659678.html, NiO-F), were used for APS and SPS processes respectively. SEM images of
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the morphology of both types of powder particles used in this investigation are shown in Fig. 1.
Figure 1: Morphology of precursor powders: (a) Ni powder and (b) NiO powder.
Suspension Preparation
For SPS process, a series of suspensions containing 10 wt% of NiO powder in ethanol as the solvent were prepared. To increase the time-stability of the particle dispersion and to prevent agglomeration of the particles, polyvinylpyrrolidone (PVP) was added as the dispersing agent. 1, 2, 3, 4 and 5 wt% of PVP were added to the suspension at different test tubes to adjust the mass percentage of the dispersant by sedimentation tests. After mixing, suspensions were sonicated for 10 minutes at a power of 60 W to break agglomerates. Figure 2 depicts sedimentation results achieved for different PVP concentrations after 14 days. It can be seen that the minimum sedimentation was obtained for the suspension containing 1 wt% PVP.
Figure 2: Sedimentation test results for 10 wt% NiO suspension and 0 to 5 wt% PVP after 14 days (freely drawn tend line).
The viscosity of this suspension was measured by a Cannon-Fenske reverse flow capillary viscometer (VWR) with an error
margin of ±0.38%. Kinematic viscosity of the suspension was measured 1.69 × 10-6 (m 2/s) at 22 oC, showing a 15.6% increase with respect to the viscosity of ethanol.
Injection Set-Up
The injection system for SPS was composed of two tanks, one containing the suspension and the other one containing ethanol. For the spraying process, the tanks were pressurized with N 2 gas. To avoid contamination and clogging of the injection system, the ethanol stored in the tank was used to clean the injection system after each suspension projection. The position of the injector, having a 200 μm diameter nozzle, was adjusted in such a way that the suspension drops penetrate adequately to the plasma jet at the nozzle exit.
Spraying Processes
Plasma spraying experiments were carried out using a Sulzer Metco 3MB atmospheric plasma-spraying gun, mounted on a computer-controlled robotic arm. Coatings were deposited on grit blasted, rectangular low carbon steel coupons of dimensions 25 × 25 × 3 mm 3. Argon was used as the primary plasma gas as well as the powder carrier gas, and hydrogen was used as the secondary gas.
To determine the optimum operating conditions for producing coatings with a high surface area, while having a minimum number of experimental runs Taguchi statistical method (Ref 10) was employed in design of the experiments (DOE). Using such an approach greatly reduce the experimental time and cost. DOE is further useful in evaluation of the importance of each selected coating variable on the surface areas of the coatings.
Table 1 presents the four variables and the three selected levels of each variable (L 9 Orthogonal array), and Table 2 shows the spraying parameters used in the APS coating processes. A fixed powder-feeding rate of 30 g/min was used, and 15 over-layers were deposited to create each APS coating. Table 3 and Table 4 show the five variables and their two selected levels of each variable (L 8 Orthogonal array) and the spraying parameters for the SPS processes. Each SPS coating was composed of deposition of 20 over-layers.
Characterization and Analysis
One of the key parameters to assess the performance of an electrode is its electrochemical active surface area. It is well known that the electrochemical double layer capacitance (ECDLC) of an electrode is proportional to the electrochemical active surface area of the material used (typically 20 μF/cm 2) (Ref 11). Therefore, for initial comparison of the electrochemical active surface area of the coatings formed by APS and SPS processes, their ECDLC were measured. A three-electrode configuration was used for all measurements. A Pt wire and an Hg/HgO electrode (saturated 1M KOH) were used as counter and reference electrodes in a 0.5M NaOH solution.
A geometrical surface
a
b
area of 0.78 cm2of each coating was exposed to the electrolyte as the working electrode. The electrolyte solution was bubbled with nitrogen during the measurements.
Cyclic voltammetry measurements were performed at five scan rates of 0.02, 0.05, 0.1, 0.15 and 0.2 V/s in the potential window from -1.35 to 0.7 V vs. Hg/HgO using an in-house potentiostat. The region around -0.1 to 0.4 V vs. Hg/HgO is considered to be essentially free of faradic current and was used to evaluate the contribution of the ECDLC to the recorded current (Fig. 3). The capacitance values are calculated by dividing the average current at the center of the positive and negative sweeps of the resulting cyclic voltammograms by the scan rates (Ref 11). Further surface measurements are needed to confirm the results.
Table 1: Variables and their levels in Taguchi design of experiments for APS.
Variables Levels (1, 2, 3) A: Standoff distance (cm) 12, 17, 22
B: Current (A) 400, 450, 500
C: Plasma gas flow (Ar/ H2)
(NLPM) 35/2.2, 35/4.4,
35/6.6
D: Torch traverse speed (m/s) 0.5, 1 Table 2: Plasma spraying parameters used for APS coatings.
Run order Standoff
distance
(cm)
Current
(A)
Plasma gas
flow
(Ar/H2)
(NLPM)
Torch
traverse
speed
(m/s)
A1 12 400 35/2.2 0.5
A2 12 450 35/4.4 1
A3 12 500 35/6.6 1
A4 17 400 35/4.4 0.5
A5 17 450 35/6.6 0.5
A6 17 500 35/2.2 1
A7 22 400 35/6.6 1
A8 22 450 35/2.2 0.5
A9 22 500 35/4.4 0.5 Table 3: Variables and their levels in Taguchi design of experiments for SPS.
Variables Levels (1, 2) A: Standoff distance (cm) 4, 6 B: Suspension flow rate (g/min) 31.2, 22.8
C: Plasma gas flow (Ar/H2) (NLPM) 50/5, 50/3
D: Current (A) 500, 450 E: Torch traverse speed (m/s) 0.5, 1 Table 4: Plasma spraying parameters used for SPS coatings. Run
order
Standoff
distance
(cm)
Suspension
flow rate
(g/min)
Plasma
gas flow
(Ar/H2)
(NLPM)
Current
(A)
Torch
traverse
speed
(m/s)
S1 4 31.2 50/5 500 0.5
S2 4 31.2 50/5 450 1
S3 4 22.8 50/3 500 0.5
S4 4 22.8 50/3 450 1
S5 6 31.2 50/3 500 1
S6 6 31.2 50/3 450 0.5
S7 6 22.8 50/5 500 1
S8 6 22.8 50/5 450 0.5
Figure 3: Cyclic voltammetry of S6 sample, and non-faradic region used for ECDLC measurements between the dashed lines (scan rate of 0.02 V/s).
Signal-to-noise (S/N) ratio, a measure of robustness, was used to analyze the influence of each spraying control factor on the specific surface area of the coatings. Since the goal is to maximize the specific surface area, the S/N ratio of “larger is better” was used. This category is calculated as a logarithmic transformation of loss function:
/ =(?10 × / ) (Eq 1)
Where n is the number of observations, and y is the observed data.
Morphology of powder particles and coating surfaces were studied by a Hitachi S-3400N VP scanning electron microscope (SEM).
Results and Discussion
Figure 4 shows the calculated double layer capacitances for all APS and SPS samples. As illustrated in this figure, by coating the sample using either APS or SPS, the ECDLC is increased significantly compared to the uncoated substrate.
Region for
ECDLC
Measurement
Nevertheless, the samples coated with the suspension exhibited a higher ECDLC compared to the APS deposited ones. As indicated above, the double layer capacitance of a sample is in a direct relation with its electrochemical active surface area. Therefore, it can be concluded that the coatings produced by nano-sized powder have a larger electrochemical active surface area in comparison with those produced by micron-sized powder.
Figure 4: Electric double layer capacitance of APS and SPS coated samples.
APS Coatings
ECDLC results, illustrated in Fig. 4, indicate that A1 with 145 μF has the highest and A7 with approximately 101 μF has the lowest capacitance among APS coated samples. Surface SEM micrographs of these two samples can be seen in Fig. 5.
Figure 5: SEM surface images of APS samples A1 and A7. The images show that the top surface of A7 coating is composed of larger splats with diameters of >100 μm, while A1 has a finer and rougher surface structure. The presence of semi-molten and re-solidified particles is evident in the latter coating. Plasma sprayed coatings are formed by deposition of flattened splats created by impact and rapid solidification of molten and semi molten particles. The size, morphology and bonding of the splats along with porosity determine the structural properties of that deposit. The different surface structures can be directly related to the plasma spraying parameters that were used to form the coatings (Ref 12). Table 5 presents the S/N response data for the electric double layer capacitance. Based on the value differences between the larger and smaller levels, the relative importance of each factor is quantified as factor ranks. The results indicate that standoff distance has the greatest influence on the electro-active surface area, followed by torch traverse speed, current and hydrogen gas flow. Figure 6 shows the effect of each control factor level on the active surface area.
Table 5: Response Table for Signal to Noise Ratios for APS. Level Standoff
distance
Current Hydrogen
gas flow
Torch traverse
speed
1 42.27 42.16 41.81 42.39
2 42.36 41.58 42.16 41.24
3 41.00 41.90 41.66 -
Delta 1.36 0.59 0.50 1.15 Rank 1 3 4 2
At 12 cm standoff distance, higher velocity and temperature of the impacting particles provide a better cohesion between the depositing layer and the substrate. Meanwhile, the smaller particles have a tendency to re-solidify because of the more rapid heat loss and deceleration at longer inflight paths. This combination of completely melted, semi-melted and re-solidified particles leads to the formation of deposits with a higher degree of porosity, rougher surfaces, and consequently larger surface areas. Increasing the standoff distance from 12 to 17 cm did not have any significant effect on the electrochemical active surface area of the coatings. By further increase of the standoff distance to 22 cm, the particle velocity and temperature are further reduced at the time of impact to the substrate. Therefore, more inflight particles at longer spray distances are re-solidified, which also causes a weaker adhesion of the particles to the substrate. In this case, only a limited number of particles, which are still molten and semi-molten by the time of impact, would adhere to the substrate and fabricate coatings consisting of larger splats with a more limited surface area. For a better understanding of the behavior of inflight particles, their temperature and velocity at the point of impact need to be determined using a DPV-2000 on-line diagnostic system.
The second parameter that affected the active surface area is torch traverse speed. The results indicate that by increasing the torch traverse speed from 0.5 m/s to 1m/s the active surface area was reduced. The authors could not find an explanation for this behavior, and further investigations are required. The results in Fig. 6 and Table 5 show that the current and hydrogen gas flow did not have a significant effect on the electrochemical active surface area.
Since the goal here is to maximize the specific surface area, the factor levels that produce the highest mean should be employed. Analysis of the results leads to the conclusion that a
A1 A7
100 μm 100 μm