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VOLUME 121 NO. 4 APRIL 2021
Contents President’s Corner: Reminiscing on the Purpose of Mining Associations by V.G. Duke . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . iv–v
PROFESSIONAL TECHNICAL AND SCIENTIFIC PAPERS Characterization of additively manufactured AlSi10Mg cubes with different porosities C. Taute, H. Möller, A. du Plessis, M. Tshibalanganda, and M. Leary . . . . . . . . . . . . . . . 143 Additive manufacturing can be used to produce components with complex and custom geometries. However, defects such as porosity and surface roughness can severely limit industrial applications. In this paper we report on the characterization of porosity, by X-ray tomography, in samples of AlSi10Mg produced by laser powderbed fusion, with porosity differences induced by changes in the process parameters. A better understanding of the defects can lead to improved design and manufacturing of additively manufactured parts. The crack growth resistance behaviour of aluminium alloy 2024-T3 at slow strain rates after exposure to standard corrosive environments C.C. Pretorius, R.J. Mostert, and S. Ramjee. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 151 The study investigates the effect of prior corrosive exposure on the crack growth resistance behaviour of thin sheet (3 mm thick) aluminium alloy 2024-T3 at slow strain rates. Compact tension specimens were exposed to standard corrosive environments that simulate accelerated atmospheric corrosion attack. It is postulated that the observed degradation of the Kc_e values of the exfoliation corrosion (EXCO)exposed material is due to hydrogen embrittlement. The specimens exposed to sodium chloride underwent a similar degradation. It is unclear what were the relative contributions of the notch effects and hydrogen embrittlement to the degradation of the KR performance.
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PROFESSIONAL TECHNICAL AND SCIENTIFIC PAPERS The use of 4-methylbenzenesulfonate ionic liquid derivatives as environmentally friendly corrosion inhibitors for mild steel in hydrochloric acid T. Nesane, S.S. Mnyakeni-Moleele, and L.C. Murulana. . . . . . . . . . . . . . . . . . . . . . . . . . . 159 Two synthesized ionic liquids were evaluated for their inhibiting effect on the corrosion of mild steel in acidic solution. Organic moieties were investigated using Fourier transform infrared spectroscopy (FTIR). Gravimetric analysis showed an increase in the inhibition efficiency as the concentration of inhibitor was increased. Polarization curves revealed that the ionic liquids act as mixed-type corrosion inhibitors. Deformation and fracture behaviour of the g-TiAl based intermetallic alloys M.N. Mathabathe, A.S. Bolokang, G. Govender, C.W. Siyasiya, and R.J. Mostert. . . . . . . . 169 Samples of b-solidifying g-TiAl intermetallic alloys of different nominal composition were produced. Microstructural examination indicated that the alloys were comprised of lamellar structures (a2+g) embedded in columnar dendritic cores in the as-cast condition. The samples showed no obvious ductility during tensile deformation, indicating the brittleness of the alloys. The fracture mode results revealed that the alloys failed by translamellar fracture with correspondingly few cleavage facets. Evaluation of spheroidized tungsten carbide powder produced by induction plasma melting R.M. Dire, H. Bissett, D. Delport, and K. Premlall. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 175 Spheroidization is a heat treatment modification process that is used to convert irregular particles to spheroidal shapes. Tungsten carbide (WC) was spheroidized using an inductively coupled radio frequency plasma at various plate powers between 9 and 15 kW. The influence of additional H2 in the sheath gas on the chemical composition of the WC was also investigated. Optical analysis of SEM micrographs indicated that the spheroidization ratio increased with increasing plasma energy.
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President’s Corner
Reminiscing the Purpose of Reminiscing onon the Purpose Mining Associations of Mining Associations ’Genius is the gold in the mine, talent is the miner who works and brings it out’ – Marguerite Gardiner
T
he concept of mining associations started in the late 1800s when the mining industry first became popularized. It was then that organizations such as the AusIMM in Australia (Australasian Institute for Mining and Metallurgy), CIM in Canada (Canadian Institute of Mining, Metallurgy and Petroleum), the IOM3 in the UK (Institute of Materials, Minerals and Mining), the AIME (American Institute of Mining, Metallurgical and Petroleum Engineers) and us, the SAIMM (Southern African Institute of Mining and Metallurgy), emerged.
The AusIMM was initially established in 1893 as the Australasian Institute of Mining Engineers after the AIME in America proved to be a major success. They at first limited membership to only industry professionals related to metallurgy, engineering, and geology. The association now admits as members and represents all individuals in the mining industry. The AusIMM focuses on the environmental, social, as well as economic aspects of the industry with the intention of improving these respective outcomes. The AIME was created in 1871 and was one of the first national engineering societies in America. The AIME has gone on to establish four more affiliated societies, namely the Society for Mining, Metallurgy and Exploration (SME), the Minerals, Metals and Materials Society (TMS), the Association for Iron and Steel Technology (AIST), and the Society of Petroleum Engineers (SPE). All these associations serve to further scientific and technological developments, both locally and internationally, to advance their respective industries. The CIM was established in 1898 as the Canadian Mining Institute. Professionals in the industry initially sought a method of influencing and communicating laws regarding worker safety, and created the CIM to achieve that. The CIM soon grew in popularity, gaining members from across Canada, and today is Canada’s leading mining association. Established in 1869, the Iron and Steel Institute held bi-annual meetings to discuss scientific questions about the manufacturing of iron and steel. As the organization grew to encompass a more general industry, it ultimately became the IOM3, which is concerned with scientific development throughout various mining operations. Its major concern is currently environmental, and the Institute is involved in dealing with climate change, greenhouse gas emissions, and general societal well-being regarding mining activities. The SAIMM was established in 1894 as a learned society which focused on academic disciplines relevant to the mining industry. It has upheld that motive throughout its existence, and still aims to spread information about scientific and technological advances that affect mining processes. After more than one hundred years, these associations have continued to grow in popularity. Many mining, metallurgical and engineering associations have been created with the intention to contribute to their industries. But what purpose do they serve to the members who form part of these associations? For one, most mining, metallurgical, and engineering associations act as academic portals, thereby providing all the information that their members need. Some even serve to further student aspirations with bursaries, scholarships, and competitions.
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Reminiscing on the Purpose of Reminiscing onAssociations the Purpose Mining of Mining Associations (continued)
Another reason why associations continue to garner support is the networking and employment opportunities that they present. Most associations highlight the developing of professional profiles of individuals and companies in the industry, thereby increasing awareness of their capabilities and service offerings. Furthermore, associations maintain a specific code of ethics and quality control in the industry. This allows them to advocate professional standards in mining and ensure that individuals uphold the standards required by industry and government legislation. Of course, there are also the rewarding aspects related to associations such as the SAIMM. These entail officially recognizing their members’ various achievements in the mining, metallurgy, and engineering industries. The associations also present awards for these achievements. Mining associations were initially formed as a connection between industry workers and the governments under which they operated, which allowed for a safe space to explore the industry and allow it to grow. This is why they are still relevant today, and why they continue to attract members. We encourage membership of these associations, if only for the extraction of one of life’s rarest commodities: real contact time. In this digital age where we obtain information and conduct meetings via electronic resources, we tend to forget how necessary in-person conversations are, and how imperative they have become. By hosting seminars and events, associations bring individuals in the mining industry together to remind us of the value that meeting face-to-face with our peers will always hold.
V.G. Duke President, SAIMM
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Characterization of additively manufactured AlSi10Mg cubes with different porosities C. Taute1, H. Möller1, A. du Plessis2,3, M. Tshibalanganda2, and M. Leary4 Affiliation: 1 Department of Materials Science and Metallurgical Engineering, University of Pretoria, Pretoria 0002, South Africa. 2 Research group 3DInnovation, Stellenbosch University, Stellenbosch 7602, South Africa. 3 Department of Mechanical Engineering, Nelson Mandela University, Port Elizabeth 6001, South Africa. 4 RMIT Centre for Additive Manufacturing, RMIT University, Melbourne 3000, Australia. Correspondence to: Anton du Plessis
Email:
anton2@sun.ac.za
Dates:
Received: 27 Aug. 2020 Revised: 12 Apr. 2021 Accepted: 19 Apr. 2021 Published: April 2021
How to cite:
Taute, C., Möller, H., du Plessis, A., Tshibalanganda, M., and. Leary, M. 2021 Characterization of additively manufactured AlSi10Mg cubes with different porosities. Journal of the Southern African Institute of Mining and Metallurgy, vol. 121, no. 4, pp. 143–150. DOI ID: http://dx.doi.org/10.17159/24119717/1331/2021 ORCID C. Taute https://orchid.org/0000-00015771-3915 H. Möller https://orcid.org/0000-00016075-9965 A. du Plessis https://orchid.org/0000-00024370-8661 M. Tshibalanganda https://orcid.org/0000-00031933-5698 M. Leary https://orcid.org/0000-00022135-1681
Synopsis Additive manufacturing can be used to produce complex and custom geometries, consolidating different parts into one, which in turn reduces the required number of assemblies and allows distributed manufacturing with short lead times. Defects, such as porosity and surface roughness, associated with parts manufactured by laser powder bed fusion, can severely limit industrial application. The effect these defects have on corrosion and hence long-term structural integrity must also be taken into consideration. The aim of this paper is to report on the characterization of porosity in samples produced by laser powder bed fusion, with the differences in porosity induced by changes in the process parameters. The alloy used in this investigation is AlSi10Mg, which is widely used in the aerospace and automotive industries. The sample characteristics, obtained by X-ray tomography, are reported. The design and production of additively manufactured parts can be improved when these defects are better understood. Keywords additive manufacturing, L-PBF, AlSi10Mg, porosity, surface roughness, density.
Introduction Additive manufacturing (AM) is fast becoming an important production method in the fourth industrial revolution, due to the possibilities it presents in terms of complex as well as custom geometries (DebRoy et al., 2018; Dilberoglu et al., 2017; Korpela et al., 2020; Tofail et al., 2018). This allows shorter lead times through reduction of parts required for assembly by merging parts (DebRoy et al., 2018; Korpela et al., 2020; Tofail et al., 2018). AlSi10Mg is popular in conventional casting methods, and substantial research effort has been applied to manufacture it successfully by AM. The addition of magnesium (Mg) gives an advantage by improving heat-treated strength due to the formation of Mg2Si precipitates (Sercombe and Li, 2016). There are three main pore types associated with AM, namely spherical pores, lack-of-fusion defects, and keyhole pores. Spherical pores are usually a result of gas that becomes trapped in the melt pool during the rapid solidification which is characteristic of laser powder bed fusion (L-PBF). They are generally very small in size. Lack-of-fusion pores are created when there is insufficient overlap in layers during the melting process. Insufficient overlap can mean that an area of poor bonding is created or, for extreme cases, unmelted powder is trapped in the remaining cavities. As overlaps are difficult to fully re-melt, lack-of-fusion pores are formed (Zhang, Li, and Bai, 2017). Keyhole pores occur in a vapourfilled depression well which collapses and forms large, rounded pores. Porosity in AM, and especially L-PBF, is influenced by laser power, scanning speed, hatch spacing, layer thickness, and energy density (Tang, Pistorius, and Beuth, 2017). When scanning speed, layer thickness, and hatch spacing are kept constant, higher laser power (and higher temperature) is expected to create deeper melt pools, which cause keyhole pore formation (Bayat et al., 2019; Khairallah et al., 2016; Mohr et al., 2020; Shrestha et al., 2019; Stugelmayer, 2018; Zhao et al., 2020). Lower laser powers are expected to lead to lack-of-fusion pore formation (Bayat et al., 2019; Majumdar et al., 2019; Mohr et al., 2020; Stugelmayer, 2018). Similar to lower laser power, faster scan speed at fixed other
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Characterization of additively manufactured AlSi10Mg cubes with different porosities parameters also leads to lack of fusion and similarly, slower scan speed leads to more keyhole pore formation. This is shown schematically in Figure 1. It has also been shown that parts never truly reach full density, i.e. 0% porosity, even at optimal process parameters. An optimal combination of processing parameters can minimize porosity to below 0.01%. The transition of lack-offusion pores is seen to be much sharper than for keyhole pores, where the transition is more gradual, as laser power is increased (du Plessis, 2019). This is illustrated in Figure 2. The optimal power for minimal porosity is seen to be lower for the lower scan speed. The alloy used in this experiment was L-PBF Ti6Al4V. Characterizing porosity is especially important, as various studies have shown it to have a detrimental effect on the mechanical properties of AM parts. This specifically includes fatigue strength, where research revealed that pores act as crack initiators and that near-surface pores are the most critical (Zerbst et al., 2019a, 2019b, 2019c). Another study of AlSi10Mg formed by AM found that areas with significant unmelted powder will undergo local cracking (Read et al., 2015). Investigations into defect formation and anisotropic properties indicated that the anisotropy of both tensile ductility and fatigue properties is intensified by defects, specifically irregularly shaped porosity such as lack-of-fusion defects (Zhang, Li, and Bai, 2017; Tang and Pistorius, 2017; du Plessis, Yadroitsava, and Yadroitsev, 2020).
Figure 1—A typical trend of part porosity with changes in scan speed and energy density (at constant power) (Tang, Pistorius, and Beuth, 2017)
Non-destructive testing (NDT) is advantageous for understanding sample integrity or density without destroying the sample. Common NDT methods are the Archimedes method, gas pycnometry, ultrasonic testing, and X-ray computed tomography (CT) scanning. The Archimedes method is relatively simple, cheap, and fast. It calculates density based on the part’s mass measured in air and in liquid (such as water or acetone). The density of the part is calculated according to Equation [1]: [1] where ρ is the part density, ρL is the temperature-dependent density of the liquid, ma is the part mass in air, and mL is the part mass in the liquid. Acetone is recommended only in its pure form, as it is hygroscopic, otherwise de-ionized or distilled water is preferred to minimize air bubbles (Spierings, Schneider, and Eggenberger, 2011). The disadvantage of this method is that it can only determine bulk density relative to the fluid used for measurement, as well as assuming 100% material density. Porosity present in the part is then determined by comparing the Archimedes density to the reference density for the material. This also means localized pores cannot be individually evaluated (Wits et al., 2016). Gas pycnometry is a process that measures part volume by displacement of an inert gas such as helium (He). Part density is calculated by measuring the mass and volume of the parts separately. As with the Archimedes method, pycnometry is relatively easy, but the disadvantages of this method are higher equipment costs and volume detection is limited only to parts that are relatively small. This method measures skeletal density, which means that the gas penetrates all open (surface connected) pores and hence excludes them from the measurement. This means that porosity is then again determined by comparing the calculated density to the reference density (Wits et al., 2016). X-ray CT scanning can be used for both dimensional and porosity analysis, among other applications (du Plessis and le Roux, 2018; du Plessis et al., 2018a). X-rays are used to form a ‘shadow’ image of the sample, as the rays are projected around and through the sample. The sample is rotated in front of a stationary X-ray source and the scan records these images from the various angles presented. This is followed by a software algorithm that calculates X-ray density at each point using backprojection, creating the 3D volume data (du Plessis, Yadroitsava
Figure 2—Porosity values as a function of laser power for two scan speeds (in mm/s) (du Plessis, 2019)
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Characterization of additively manufactured AlSi10Mg cubes with different porosities and Yadroitsev, 2020). The final result is a high-quality 3D image with a clear view of where pores are located, as well as the size range of the pores (du Plessis et al., 2018a). A comparison of NDT methods by Wits et al. (2016) indicated that CT scanning can measure pore areas that are smaller than the detectable size for microscopic methods, as well as predict densities more accurately than the Archimedes method. It also allows further analysis of the porosity present in the parts, such as sphericity, pore distribution, and defect volume both on the surface and inside the sample (du Plessis, Yadroitsava and Yadroitsev, 2020; Wits et al., 2016).
Materials and methods This study used AlSi10Mg samples that were printed at RMIT University in Melbourne, Australia. A total of 25 solid cubes were printed using an SLM500 Quad laser system from SLM Solutions, with dimensions 10 × 10 × 10 mm3, and material composition as described in Maconachie et al. (2020). Differences in porosity content were induced by varying the printing process parameters, specifically the laser power. Five different power settings were used, namely 210, 280, 350, 420, and 490 W, with five samples printed for each power setting. The other parameters used in the printing are listed in Table I. Different characterization methods were used as described below. X-ray CT scanning was used in two ways: once to measure the volume of the cube for a density analysis method based on volume and scale mass, and once to do a conventional CT-based porosity analysis. The system used was a GE Nanotom S and the software used for image analysis was Volume Graphics VGSTUDIO MAX 3.3.
Initial testing Initial testing done at the University of Pretoria included measuring dimensional accuracy and mass, using a New Classic ML Mettler Toledo scale, which has a draft shield to improve accuracy. The mass in air was taken as an average of three to four individual measurements. These values were used in the subsequent density calculations.
Density measurements Three methods were used to determine the bulk density of the samples (average density), namely Archimedes, gas pycnometry, and a CT-based density method described previously in du Plessis et al. (2018b). Archimedes density measurements were done using a New Classic ML Mettler Toledo scale, accurate to four decimals, and distilled water. Four individual measurements of the samples in water were taken. The water temperature was measured to ensure accurate water density was used. Density was then calculated according to Equation [1]. Gas pycnometry was carried out using a Micromeritics AccuPyc II 1340 gas pycnometer with helium gas at a calibrated pressure of approximately 19.5 psig (134.45 kPa). Five volume measurements were obtained per sample, and the average used to calculate density. The CT-based density calculations were based on the mass scale of the samples in air and the CT-determined volume of the cube, segmented carefully to include all pore spaces.
NanoCT scan porosity measurement The analysis method used a procedure identical to that outlined in du Plessis et al. (2018c), which minimized bias in the The Journal of the Southern African Institute of Mining and Metallurgy
segmentation process. Porosity percentage values are used here, despite much more information being available. This additional data will be used in future work for further detailed analysis of pore morphologies.
Surface roughness Surface roughnesses of the samples were obtained using optical microscopy with an Olympus DSX 510 at Wirsam Scientific. The microscope uses Olympus Stream software to plot a surface map of the sample and return surface roughness values, such as arithmetical mean height (Sa) along with a colour-scale map of the surface. Surface roughness was measured in a minimum of three separate areas (1960 µm × 1960 µm each) on both the top surface and the side surfaces, to obtain an overall representative average. The distance between each layer scanned by the microscope was 12.1 µm, whereas the height range scanned differed between samples due to different surface conditions.
Results and discussion Three density methods were used to calculate the difference in density of the samples from the different laser power sets. This is shown in Table II. An increase in laser power is shown to lower the measured bulk density of the samples, and all three methods are consistent in this trend. In the table A refers to Archimedes density, GP is gas pycnometry density, and CT is the CT scan density. AlSi10Mg has a theoretical density of approximately 2.68 g/cm³. From Figure 3 it can be seen that the pycnometer and CTbased densities correlate well, whereas the Archimedes density is lower. This lower value might be attributed to air bubbles attached to the surface of the sample when submerged in water, which affects the measured mass of the sample in water. The rough surface is conducive to air bubbles attaching and this was physically observed. CT scan images of one representative sample from each laser power set are shown in Figure 4. From left to right the power settings were 210, 280, 350, 420, and 490 W. It can be seen that the first two cross-sections on the left have very small amounts of porosity, which corresponds to the lower power settings. The Table I
Printing parameters used Parameter
Layer thickness (µm) Laser velocity (mm/s) Hatch spacing (µm) Scan strategy
Value 50 921 190 One contour scan followed by hatch tracks in zigzag pattern with 90 degree change per layer
Table II
ensity of samples of each laser power set, according D to three measurement methods Laser Power (W)
210 280 350 420 490
2.550 2.549 2.512 2.431 2.365 VOLUME 121
Density (g/cm³)
A GP CT 2.610 2.608 2.574 2.495 2.419 APRIL 2021
2.613 2.609 2.576 2.516 2.444
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Figure 3—Density calculated with three different methods, as a function of laser power, with error bars
Figure 4—Representative CT cross-sectional images for each of the five different laser power settings, showing porosity (black dots) and also indicating the presence of some dense inclusions (white dots encircled in red). Arrow indicates build direction
last two cross-sections on the right are seen to contain a greater number of pores, which are larger and are rounded, which corresponds to the higher laser power and subsequently higher energy density. This large, rounded porosity at high power is attributed to keyhole mode porosity formation. The encircled areas in Figure 4 show the presence of high-density inclusions, which could potentially be due to contamination from the powder itself, as Al alloys usually contain iron (Fe) impurities. The system used by RMIT University does not print other metal alloys, so the contamination is not from a previous build. Higher density particles from the powder itself could include iron (approx. 7.9 g/cm³), chrome (approx. 7.2 g/cm³), manganese (approx. 7.3 g/cm³), nickel (approx. 8.9 g/cm³), titanium (approx. 4.5 g/cm³), or copper (approx. 8.9 g/cm³). It is also quite possible that some of the inclusions are Al2O3 (3.99 g/cm³) if oxidation occurred during spattering, even in the inert gas atmosphere used for printing. The density of the inclusions is higher than that of the alloy, thus they appear brighter in the scans. It is clear that as laser power increases, the number of high-density inclusions decreases. This can be attributed to higher temperatures or larger or deeper melt pools creating more melting/remelting and homogenization of the material. The samples have pins on the upper surface to help identify them and to keep samples from the different laser power sets separated. The number of pins indicates which laser power was used, with one pin referring to 210 W, up to five pins referring to 490 W. The green vertical arrow indicates an upwards building direction. Table III shows the total percentage of porosity, obtained from the CT results, corresponding to each laser power set, as well as
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the relationship between power and porosity in a graph insert. As seen in Figure 4 and in Table III, the lowest power (E1) sample has slightly greater porosity than E2. This can be explained as due to lack of fusion porosity at 210 W and the ideal melting with lowest porosity content at 280 W compared to 210 W. The difference is statistically significant, as the 210 W laser power resulted in an average of 0.16% porosity with a variance of 0.0024 and standard deviation of 0.0495, whereas the 280 W laser power resulted in an average of 0.08% porosity with a variance of 0.0001 and standard deviation of 0.0103. As laser power increases further, the total porosity increases, due to more keyhole porosity, which is consistent with previous work (du Plessis, 2019). The error bars are calculated using the difference between the average porosity and the maximum and minimum porosity values for each laser power set. The CT porosity analyses are shown in 3D representations in Figure 5. The porosity percentages of those specific region-ofinterest (ROI) cubes are added as inserts, with the overall average of the porosity from all 10 × 10 × 10 mm³ cubes in each laser power set in parentheses. Here it is clearly seen that the lower laser power samples have relatively little porosity, especially compared to the samples manufactured at the higher laser powers. The higher laser power samples are seen to have much higher porosity and the pores are more spherical in shape, whereas the lower power samples have less spherical, lack-of-fusion type porosity. The presence of lack-of-fusion porosity at high values in E1, compared to E2, is not consistent with the measured density values in Figure 3 and Table II and therefore requires explanation. Despite the small average values of 0.16 and 0.08%, all density The Journal of the Southern African Institute of Mining and Metallurgy
Characterization of additively manufactured AlSi10Mg cubes with different porosities Table III
Porosity percentage of each power set and plotted as a function of laser power Laser power (W)
Average porosity (%voids)
210 0.16
280
0.08
350
0.56
420
2.40
490
5.59
methods showed the E1 sample to be slightly denser. The most likely explanation for this observation is that the bulk density measurements inaccurately measure the bulk when irregularly shaped pores are present on the surface. Such pores create open cavities, allowing water or gas to enter the object in surfaceconnected pores, and are thus excluded from the measurements. The CT-based segmentation also might select more of this type of porosity or may be inaccurate – the CT volume measurement for bulk density requires accurate calibration of the voxel size, unlike the porosity percentage value in Table III and Figure 5. The shapes of the pores are clearly more irregular at the lowest laser power and more spherical at higher power. The long, irregular pores are seen clearly in E1 versus the more spherical pores at higher laser powers. A scale representation is inserted to show the lengths of each side for all five cubes. Figure 6 shows close-ups of the top surfaces of the representative samples for each laser power. These images were also obtained using VGSTUDIO MAX. As laser power is increased it can be seen that the surface roughness decreases. The first two top surfaces are seen to have a much higher surface roughness than the last two, corresponding to lower and higher laser power, respectively. The surface roughness values, Sa, are included as inserts in the figure for visual comparison. The values themselves were obtained using optical microscopy with an Olympus DSX 510, which is not dependent on track orientation relative to scanning as it is a surface area scan, which takes the Ra line profile parameter and expands it into three dimensions. The top surface was analysed as it is the final layer in the printing process, which means it is likely to be the most affected by process parameters. The bottom surface was excluded as the samples were printed on supports, leading to an irregular surface that is not representative of the parameters. The solitary pin on the samples indicates that the samples were the first in each laser power set. The pins that range from one to five indicate the laser power setting that was used, as indicated by the powers in the inserts. Figure 7 shows how the surface roughness in general decreases with an increase in laser power. The decrease is most likely due to wider and more overlapping melt pools, which creates a relatively smoother top surface. The error bars show that the surface roughness varied greatly in each laser power set. The Journal of the Southern African Institute of Mining and Metallurgy
The error bars are calculated using the difference between the average surface roughness, Sa, and the maximum and minimum Sa values for each laser power set. The side surfaces were measured to obtain a general idea of the surface roughness on the sides, to see how they differ from the top surfaces. The sides of the samples showed a much smaller variation between laser power settings, averaging between 9 and 11 µm Sa across all laser powers. This effect is shown in Figure 8. The errors bars also show how, at the highest and lowest laser powers, the surface roughness varies much more than for the middle laser powers.
Conclusions From the results it can clearly be seen that higher laser power induces a larger volume of porosity. The higher power leads to more keyhole-type porosity, whereas the lower power samples have more lack-of-fusion type pores. The results show that highdensity inclusions decrease in volume with higher laser power, due to remelting, or deeper melt pool penetration. While higher laser power seems to decrease the surface roughness, it comes at the cost of larger volumes of porosity. If only the variation of porosity with laser power is considered, the optimal power for the given scan speed is in the range 210–280 W, with the lowest porosity at 280 W. However, when considering the measured surface roughness, the lowest porosity values for the top surface are obtained at higher powers. Therefore, depending on the requirements, a suitable combination of roughness and porosity minimization can be obtained. These results indicate some of the challenges associated with L-PBF. Typically, values for porosity < 0.5% are considered reasonable and a roughness, Sa, in the range 20–24 µm might be acceptable for some applications. Understanding porosity formation and pore morphology associated with laser-powder bed fusion manufactured parts aids in improving parts to decrease the limiting effect these defects can have on parts in industry. Future work will include investigation of the effect of porosity and surface conditions on corrosion and mechanical properties. Work is also ongoing on detailed 3D pore morphology evaluation using the CT data obtained in this work. VOLUME 121
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Characterization of additively manufactured AlSi10Mg cubes with different porosities
Figure 5—Porosity analyses, for a 4 × 4 × 4 mm3 region of interest, of representative samples for each of the five laser power settings, from (top) low power to (bottom) high power, with actual porosity value of ROI as inserts and overall averages in parentheses
Acknowledgments The Light Metals Development Network (LMDN) is acknowledged for their financial support. The Collaborative Programme for Additive Manufacturing (CPAM), funded by the South African
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Department of Science and Innovation, is acknowledged for additional funding. The authors would like to thank the Stellenbosch University Research group 3DInnovation staff (Ms Carlyn Wells) for their contribution to the success of this project. The Journal of the Southern African Institute of Mining and Metallurgy
Characterization of additively manufactured AlSi10Mg cubes with different porosities
Surface roughness (Sa ), mm
Figure 6—Close-up sections of the top surface of each representative sample for each of the five laser power settings, with Sa values as inserts
Surface roughness (Sa ), mm
Figure 7—Plot of top surface roughness, Sa, as a function of laser power, with error bars
Figure 8—Plot of side surface roughness, Sa, as a function of laser power, with error bar
The authors acknowledge use of facilities in the RMIT Advanced Manufacturing Precinct Special thanks to Mr Dewald Noeth and Ms Colleen Syrett at Wirsam Scientific for the use of their Olympus DSX 510 microscope. Special thanks also to the staff at the University of Pretoria (Dr Robert Cromarty, Mr Sibusiso Mahlalela, Mr Dirk Odendaal, Mr Mfesane Tshazi) for all their assistance during this project. The Journal of the Southern African Institute of Mining and Metallurgy
Authors’ contributions Heinrich Möller: supervision. Anton du Plessis: supervision. Muofhe Tshibalanganda: nano-CT scanning. Martin Leary: printing and provision of samples. Carlien Taute: sample preparation, density and surface roughness measurements. The article was written by Carlien Taute and reviewed by all the authors. VOLUME 121
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Characterization of additively manufactured AlSi10Mg cubes with different porosities References
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The Journal of the Southern African Institute of Mining and Metallurgy
The crack growth resistance behaviour of aluminium alloy 2024-T3 at slow strain rates after exposure to standard corrosive environments Affiliation: 1 Department of Materials Science and Metallurgical Engineering, University of Pretoria, South Africa. Correspondence to: C.C. Pretorius
Email: ccep@live.co.za
Dates:
Received: 31 Aug. 2020 Revised: 25 Feb. 2021 Accepted: 13 Mar. 2021 Published: April 2021
How to cite:
Pretorius, C.C., Mostert, R.J., and Ramjee, S. 2021 The crack growth resistance behaviour of aluminium alloy 2024-T3 at slow strain rates after exposure to standard corrosive environments. Journal of the Southern African Institute of Mining and Metallurgy, vol. 121, no. 4, pp. 151–158.
C.C. Pretorius1, R.J. Mostert1, and S. Ramjee1
Synopsis The study investigates the effect of prior corrosive exposure on crack growth resistance behaviour of thin sheet (3 mm thick) aluminium alloy 2024-T3 at slow strain rates. Compact tension specimens were exposed to standard corrosive environments that simulate accelerated atmospheric corrosion attack. Two corrosive environments were considered – an exfoliation corrosion (EXCO) solution and a 3.5 wt% sodium chloride solution. The unloading compliance R-curves of the two-hour EXCO-exposed specimens revealed a significant degradation of approximately 11% in the crack growth resistance behaviour (Kc_e values) compared to the baseline (air-exposed) values. Furthermore, secondary intergranular crack formation was also revealed in the plastic zone ahead of, and adjacent to, the crack tip of these specimens; which formed during the crack growth resistance loading. It is postulated that the observed degradation of the Kc_e values of the EXCO-exposed material is due to hydrogen embrittlement since the exposure times for the EXCO evaluation were limited to ensure that uniform corrosion dominated; that is, significant penetration of corrosion damage and pitting due to localized corrosive attack did not occur. The sodium chloride-exposed specimens revealed a similar degradation (13%) after 24 hours exposure. However, slight intergranular corrosive attack and isolated pitting were observed on the exposed surfaces prior to crack growth resistance loading, resulting in notch effects that could assist in crack growth. Pitting and intergranular corrosion were, however, not observed at the pre-crack tip. The relative contributions of the notch effects and the hydrogen embrittlement during the degradation of the KR performance are, therefore, unclear. Keywords corrosion, crack growth, aluminium alloy 2024-T3.
DOI ID: http://dx.doi.org/10.17159/24119717/1340/2021 ORCID C.C. Pretorius https://orchid.org/0000-00028408-8390 R.J. Mostert https://orchid.org/0000-00028592-1313 S. Ramjee https://orchid.org/0000-00030104-207x
Introduction The long-term service usage of the popular high-strength aluminium alloy 2024 in the aeronautical industry may lead to degradation in the alloy’s mechanical properties due to atmospheric corrosion exposure. Various authors have described a decrease in the strength and ductility of the alloy during its service life (Alexopoulos et. al., 2012, 2016; Kamoutsi et. al., 2006; Sharp et. al., 1998; Sharp, Mills, and Clark, 2001). In the aeronautical industry, this degradation may be brought on by a number of embrittlement mechanisms; including the natural ageing of the alloys, exfoliation corrosion, and hydrogen embrittlement (Alexopoulos et. al., 2016). The well-established natural ageing of the alloy is a result of the growth and coarsening of the strengthening precipitates over time at ambient temperatures. Growth of the strengthening precipitates ultimately leads to a loss of coherency with the aluminium matrix; which is accompanied by the loss of the so-called coherency strains that hinder dislocation glide and provide strengthening. This may be followed by coarsening - in which the larger precipitates grow at the expense of the smaller ones (Ratke and Voorhees, 2002) - which will reduce the inter-precipitate spacing and aid in dislocation glide due to fewer obstacles hindering dislocation mobility. Moreover, owing to the general incoherency of the larger precipitates, dislocations can no longer intersect the precipitates. Accordingly, dislocations have to loop around the precipitates for dislocation motion to persist. In so doing, dislocation loops
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The crack growth resistance behaviour of aluminium alloy 2024-T3 at slow strain rates are formed around the precipitates, which act as dislocation sources that ultimately decrease the ductility of the alloy. It has been well established that the ultimate strengthening is brought on by the coherent and semi-coherent precipitates; due to the aforementioned coherency strains and relatively high precipitate density per unit area that accompany such precipitates. With the growth and coarsening of the precipitates (and ultimate incoherency), a lower mechanical strength and ductility often results. More specifically, this has a drastic effect on the proof strength (or yield strength) and the final elongation to failure; as can be seen in Figure 1 for aluminium alloy 2024-T3 artificially aged at 170°C at various ageing times (Alexopoulos et. al., 2016). Atmospheric corrosion of aluminium alloy 2024-T3 in the aeronautical industry is widely relevant due to the structural components manufactured from the alloy being regularly exposed to moisture and other corrosive environments during their service lives. With regard to the aeronautical industry, the most common manifestations of atmospheric corrosion include pitting and exfoliation. Kamoutsi et. al., (2006) showed that atmospheric corrosion of the aluminium alloy 2024-T3 is initiated by a localized breakdown of the protective surface oxide layer, which is then followed by localized corrosion (in the form of pitting) before corrosion proceeds. Similarly, Alexopoulos et al., (2012) investigated the effect of corrosion exposure time for ultra-thin sheet aluminium alloy 2024-T3, which showed similar results. The authors also showed that intergranular corrosion proceeds after the corrosion pit reaches a critical size, and grows laterally (perpendicular to the corrosion pits and parallel to the surface) in a direction parallel to the rolling direction (Alexopoulos et al., 2012). The ASTM G34-1 standard (ASTM International, 2018) defines exfoliation as corrosion that proceeds laterally along planes parallel to the surface (generally along grain boundaries), where it forms corrosion products that force the matrix apart. The standard also provides guidance for producing this degradation in an accelerated manner. Specifically pertaining to 2XXX and 7XXX aluminium alloys, the Exfoliation Corrosion (EXCO) solution was developed to simulate accelerated exfoliation corrosion attack, and is said to provide a useful prediction of the exfoliation corrosion when exposed to marine atmospheres (ASTM
International, 2018). Exposure to such corrosive environments ultimately gives rise to a layered surface appearance (ASTM International, 2018), as shown in the work by Sharp, Mills, and Clark (2001) on aluminium alloy 7075-T651 after exposure to the EXCO solution for 48 hours (refer to Figure 2). The presence of such an exfoliated surface layer will, of course, affect the mechanical properties in a deleterious way. Various authors have described a loss in both strength and ductility with increasing corrosion exposure time (Alexopoulos et. al., 2012, 2016; Kamoutsi et.al., 2006; Sharp et. al., 1998; Sharp, Mills, and Clark, 2001). Furthermore, as a result of localized stress concentrations (notch effects) at the exfoliated surface, the probability that the corrosion will interact with other damage mechanisms also increases with the extent of exfoliation (Kamoutsi et. al., 2006). An example of such interaction may be seen by referring again to the work by Sharp, Mills, and Clark (2001), which shows that the exfoliated surface resulted in the initiation of a fatigue crack in aluminium alloy 7075-T651 (Figure 2). A number of mechanical testing investigations of precorroded aluminium alloy 2024 have shown that the deleterious effects on the strength and ductility of these alloys are directly related to the corrosion exposure time (Dovletoglou et. al., 2018; Pantelakis, Daglaris, and Apostolopoulos, 2000; Petroyiannis et. al., 2004). However, these experiments also showed that, although the residual strength of the alloy can be recovered after removal of the corrosion layer, the ductility cannot (Dovletoglou et. al., 2018; Kamoutsi et. al., 2006; Pantelakis, Daglaris, and Apostolopoulos, 2000; Petroyiannis et. al., 2004). Furthermore, the studies also showed that moderate heating of the alloy after removing the corroded surface layer assists in the partial recovery of the ductility. This has been proposed as experimental evidence that the degradation of the ductility of aluminium alloy 2024 during corrosive attack is accompanied by a corrosion-induced mechanism of hydrogen embrittlement (Kamoutsi et. al., 2006). It was proposed that atomic hydrogen is introduced to the alloy during the atmospheric corrosion process through the reduction of water according to Equation [1], followed by adsorption of the hydrogen atoms on the surface (Azofeifa et. al., 1997).
Figure 1—The effect of artificial ageing at 170°C on the proof strength (engineering yield strength, σ0) and elongation to fracture (A5) of the aluminium 2024-T3 alloy ( Alexopoulos et. al., 2016)
Figure 2—Macro exfoliation of Al 7075-T651 alloy after 48 hours exposure to an EXCO solution: (a), resulting in the initiation of a fatigue crack (b) (Sharp, Mills, and Clark, 2001)
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The crack growth resistance behaviour of aluminium alloy 2024-T3 at slow strain rates The atomic hydrogen may then either be absorbed into the aluminium matrix, or react on the surface to form hydrogen gas (H2). In the former case, the atomic hydrogen will diffuse to certain preferred lattice sites such as grain boundaries, line defects, matrix/precipitate interfaces, and point defects; which act as hydrogen traps (Charitidou et. al., 1999). The presence of the hydrogen at these sites may result in embrittlement of the alloys through one, or a combination, of the conventional hydrogenembrittlement mechanisms (Charitidou et. al., 1999). These are (i) the hydrogen-enhanced decohesion embrittlement mechanism (HEDE), (ii) the hydrogen-enhanced localized plasticity (HELP) mechanism, (iii) the adsorption-induced dislocation emission (AIDE) mechanism, and (iv) hydride formation and fracture (Birnbaum, 1990; Birnbaum et. al., 1997; Lynch, 2011; Milne, Rictchie, and Karihaloo, 2003; Troiano, 1974). Although the details of these mechanisms will not be discussed here, it is important to note that they have one aspect in common; that is, hydrogen accumulates at regions of triaxial stress (or strain). In recent work by Alexopoulos et al. (2016), slow strain rate tensile tests were performed on the aluminium alloy 2024-T3 after short-term exposure to an EXCO solution. The investigation showed that the degree of embrittlement correlates with the size, location, and number of S-type (Al2CuMg) precipitates. This may provide an understanding of the role played by the triaxial elastic strains that accompany the coherent and semi-coherent precipitates during the hydrogenembrittlement process. It is obvious that a number of negative implications arise from the fact that the alloy suffers from such embrittling effects; especially considering that this may occur under atmospheric conditions. Examination of certain photographs in the published work by Alexopoulos et al. (2016) led to the impression that the application of plastic strain to EXCO-embrittled tensile specimens after short-term exposure resulted in the formation of numerous surface cracks prior to the fracture of the specimens (Figure 3). The initiation of such surface cracks at relatively low strains, or alternatively, in the vicinity of stress concentrations (such as a notch or a fatigue crack), poses a significant risk to the structural integrity of components constructed from these alloys. This is especially true in the aeronautical industry, in
which some aluminium-based structures are exposed to marine atmospheres for significant durations. In the current work, this effect is investigated using advanced analytical techniques such as X-ray microcomputed tomography (MicroCT). Furthermore, the assumed effect of EXCO-induced hydrogen embrittlement on the crack growth resistance behaviour of the alloy has not been studied extensively. The present work is, therefore, concerned with the ability of thin sheet (3.16 mm thickness) aluminium alloy 2024-T3 to resist the extension of a pre-existing (fatigueinduced) crack after being exposed to an environment that provides the necessary conditions for hydrogen embrittlement and/or corrosion.
Material and specimen preparation The compact tension specimens used for the crack growth resistance (KR) evaluation were machined from 3.2 mm thick plate material of aluminium alloy 2024-T3. As such, various specimen preparation steps were required prior to the crack growth resistance evaluation, which included the machining and fatigue pre-cracking of specimens, as well as prior exposure to the different environments considered. Preparation of the relevant exposure environments was also required prior to the mechanical testing. The following sections provide a brief description of all the preparation steps prior to crack growth resistance evaluation.
Preparation of exposure environments The KR specimens were pre-exposed to three environments: (i) air for the baseline KR values, (ii) a 3.5% sodium chloride solution, and (iii) an EXCO solution. A brief description of the preparation methodology and relevant ASTM standards for the latter two exposure environments is given.
Sodium chloride solution The 3.5% sodium chloride solution was prepared in accordance with the ASTM G44-99(2013) standard (ASTM International, 2013). After buffering with hydrochloric acid, the solution pH was measured at 6.5.
EXCO solution The Exfoliation Corrosion exposure environment is a solution comprising approximately 4.0 M sodium chloride (NaCl), 0.5 M potassium nitrate (KNO3), and 0.1 M nitric acid (HNO3). The solution was prepared in accordance with section 7 of the ASTM G34-01 standard (ASTM International, 2018). The pH of the solution was measured at 0.27.
Crack growth resistance specimen preparation Machining and fatigue precracking.
Figure 3—Scanning electron image of a peak-aged (48 hours at 170°C) aluminium alloy 2024 after 2 hours of exposure to an exfoliation corrosion solution, showing secondary cracking (white arrows) near the exposed surface (Alexopoulos et. al., 2016) The Journal of the Southern African Institute of Mining and Metallurgy
Machining commenced with the sectioning of blank squares (60 mm × 60 mm) from the 3.2 mm thick plate material of aluminium alloy 2024-T3, providing the basis for compact tension (C(T)) specimens as prescribed in section 8.3 of the ASTM E561-15a standard (ASTM International, 2019). The blank specimens were then notched to produce C(T)(L-T) type crack growth resistance test specimens, with a fatigue starter notch (radius 0.12 mm) in accordance with the ASTM E561-15a standard. The specimens were further processed by the mechanical testing laboratory at CSIR to introduce the required fatigue pre-cracks using an Instron 1342 servo-hydraulic testing machine operating on Instron crack propagation software. The specimen geometry and relevant specimen dimensions are shown in Table I and Figure 4. VOLUME 121
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Figure 4—Crack growth resistance C(T)(L-T) specimen geometry with (a) a schematic showing the relevant specimen dimensions, and (b) a photograph of the specimen and specimen notch configuration
Table I
pproximate lengths of the relevant dimensions of A the crack-growth-resistance C(T)(L-T) specimen as depicted in Figure 4 WT
W
B
an a0 afat dp
60.00 mm 48.00 mm 3.16 mm ≈ 22.50 mm ≈ 24.50 mm ≈ 2.00 mm 12.50 mm
KR specimen preparation prior to environmental exposure Further specimen preparation was required prior to the exposure of the specimens to the different environments. The preparation consisted of shielding the most of the specimen surface area from the corrosive, ensuring that only a 10 mm wide section near the specimen notch/precrack configuration would be exposed. This was achieved by initially covering the relevant areas with a layer of double-sided tape, followed by a double layer of adherent PVC tape. It should also be noted that the relevant specimen dimensions for the crack growth-resistance evaluation (i.e. W, B, and an) were measured prior to the shielding and exposure procedures.
The sodium chloride exposure procedure After the shielding of the nessasary specimen surface areas, 100 ml of the sodium chloride solution was added to a container capable of holding the 60 mm × 60 mm C(T) specimen whilst ensuring full coverage of the relevant surfaces. The volume of 100 ml provides the required minimum volume to exposed surface area ratio of 10 ml/cm2 as prescribed in section 11 of the ASTM G34-01 standard (ASTM International, 2018). The specimen was then placed in the container, and exposed to the soluton for 24 hours at room temperature. After exposure, the specimen was removed from the container and the exposed surface was lightly cleaned with acetone before removing the shielding and performing the KR evaluation.
The EXCO solution exposure procedure The exposure of the specimens to the EXCO solution was
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performed in a similar manner as for the sodium chloride solution; the only difference being a decrease in the exposure time to 2 hours.
Experimental procedure Crack growth resistance evaluation The crack growth resistance evaluations for all the relavent exposure conditions were performed using an MTS Criterion Series 45 static test system operating on MTS Elite Test Suite software. As per the unloading compliance method of the ASTM E561-15a standard (ASTM International, 2019), the software was specially programmed to introduce a number of unloading/reloading sequences during each test under crosshead extension control. Furthermore, the grips and fixtures prescribed in ASTM E1820 were used to install the C(T)(L-R) specimen onto the testing system’s loading axis. Additional steps were included during installation and testing so as to ensure maximum reproducibility in the test results. These are (i) ensuring good and similar alignment of the test specimens with the test system’s loading axis (achieved through the use of spacers during specimen installation), (ii) the application of a small (0.2 kN) preload prior to installation of the crack mouth opening displacement (CMOD) gauge, and (iii) the use of the same reference specimen for the CMOD gauge zero value. The KR testing then commenced with a crosshead extension rate of 0.03 mm∙min-1 until a crack mouth opening displacement of 3.0 mm was achieved, along with a total of 21 equally spaced (w.r.t. displacement) unload/reload sequences. Thereafter, the specimens were subjected to a final loading at a significantly higher crosshead extension rate (300 mm∙min-1) until complete separation. The stable crack extension was then measured using a stereo microscope, and taken as the average of nine measurements over the thickness of the specimen.
Scanning electron fractography Scanning electron fractography was performed using a Joel JSM IT300 scanning electron microscope (15 kV) in order to compare the fracture surface appearances under the different exposure conditions. The Journal of the Southern African Institute of Mining and Metallurgy
The crack growth resistance behaviour of aluminium alloy 2024-T3 at slow strain rates Microcomputed tomography The KR test procedure for the specimens sent for X-ray MicroCT was interrupted prior to final fracture and separation. After the KR test terminated at a 3 mm CMOD reading, the specimens were sent for EDM in order to produce a smaller specimen that contained the notch/crack configuration. The specimens were then subjected to MicroCT scanning at the University of Stellenbosch.
Results and discussion Baseline crack growth resistance results The baseline (air-exposed) crack growth resistance curves are shown in Figure 5. In accordance with the compliance method prescribed in the ASTM E561-15a standard, an average apparent resistance to crack extension (KC_e) of 91 ± 2.8 MPa√m was obtained at an effective crack extension (∆ae) of 6.14 ± 0.43 mm. This value is significantly higher than that reported (61.6 ± 8.0 MPa√m) by Reynolds (1996). The deviation is believed to be due to (i) the low displacement rate used in the current investigation, and (ii) the reduced size of the remaining ligament of material – both of which will alter the degree of plasticity ahead of the crack front. With regard to the former, owing to the low displacement rate, the initial rate of change in stress intensity does not fall within the limits of 0.55–2.75 MPa√m∙s-1 (as set out by the ASTM E561-15a standard). The reduced impact of the loading rate increases the probability for crack tip blunting, since a larger fraction of the material (ahead of the crack) may absorb the applied energy through plastic yielding. The reduced size of the remaining ligament of material will have a similar effect on C(T)-type specimens, with an increase in the plastic zone size and a larger fraction of material absorbing the applied energy. Consequently, the values determined in this study do not represent the plane-stress fracture toughness of alloy 2024-T3 at a thickness of 3.2 mm. It is for these reasons that the normal nomenclature, as given in the relevant ASTM standards, is not used to describe the crack growth resistance values. Nevertheless, due to the comparative nature of the investigation, the values may be used to establish whether the exposure environments alter the crack growth response of the alloy. In this respect, the increase in plasticity in point of fact benefits the investigation, since embrittling effects should alter the plastic response through either the hindrance or localization of plastic flow.
Environmentally exposed crack growth resistance results The KR curves for the environmentally exposed and baseline results are compared in Figure 6. A significant degradation may be seen in the crack growth-response of both the sodium chloride- and EXCO-exposed materials. With regard to the former, a reduction in the apparent crack growth resistance (KC_e) of approximately 13.0 ± 0.5% is observed, with a KC_e of 79.0 ± 0.5 MPa√m. Similarly, the EXCO-exposed material revealed a reduction of approximately 11.2 ± 0.1% with a KC_e of 81.0 ± 0.7 MPa√m. The degradation in the crack growth resistance behaviour is recorded in more detail in Table II.
Fracture surface appearance and MicroCT mapping: The fracture surface appearance of the EXCO-exposed material is shown in Figure 7. From this fractograph, it is clear that the fracture surface may be subdivided into four regions. That is, (A) the fatigue precrack perpendicular to the applied load, (B) the The Journal of the Southern African Institute of Mining and Metallurgy
Figure 5—The apparent crack growth resistance behaviour of 3.2 mm thick aluminium alloy 2024-T3 after prior exposure to air (baseline results)
Figure 6—Comparison between the crack growth resistance behaviour after exposure to different corrosive environments. Note the degradation in the apparent crack growth resistance (KC_e) after prior exposure to the EXCO and NaCl environments
stable crack extension fracture surface (also perpendicular to the applied load), (C) the shear (ductile) type fracture surface with an orientation of approximately 45° to the applied load, and (D) an intergranular type fracture surface near the edge of the specimen. The fracture surface types (A) to (C) were found to repeat during all the exposure conditions investigated. However, only the EXCO-exposed specimens revealed the type (D) fracture surface appearance. The high-resolution scanning electron fractograph in Figure 8 shows the transition from intergranular (type D) to dimple fracture for the EXCO-exposed material in more detail, with Figure 9 showing higher magnification scanning electron fractographs of the fracture surfaces labelled C (Figure 9i) and D (Figure 9ii). It becomes apparent that significant weakening of the grain boundary strength has occurred during the two-hour exposure to the EXCO solution. It is also found that the intergranular fracture surfaces correlate well with the VOLUME 121
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The crack growth resistance behaviour of aluminium alloy 2024-T3 at slow strain rates Table II
ummary of the apparent crack-growth resistance values of the baseline (air-exposed) and exposed aluminium alloy S 2024-T3, showing the percentage degradation as compared to the baseline values Specimen label 35 36 Average (Air)
T3, as received T3, as received –
Specimen label 6 7 Average (NaCl)
2
Condition
Condition T3, 24h NaCl exposure T3, 24h NaCl exposure –
Specimen label
Condition
2 3 Average (EXCO)
T3, 2h EXCO-exposure T3, 2h EXCO-exposure –
Baseline (air exposed) material KC_e (MPa√m) 89.0 93.0 91.0 ± 2.8
Sodium chloride-exposed material KC_e (MPa√m) 79.5 78.9 79.2 ± 0.4
EXCO-exposed material KC_e (MPa√m) 81.3 80.3 80.8 ± 0.7
Degradation (%) Zero Zero Zero
Degradation2 (%) 12.7 13.3 13.0 ± 0.5
Degradation (%) 10.7 11.7 11.2 ± 0.1
Calculated as (Kc(average baseline)-Kc(i, exposed)) / Kc(average baseline)
Figure 7—Post-crack growth resistance fractograph of the Exco-exposed material, showing the positions of the different fracture surface types; (A) fatigue precrack, (B) the stable crack extension, (C) shear fracture opposite the stable crack extension, and (D) the intergranular fracture surface
secondary crack formation (refer to the MicroCT scan in Figure 11) that resulted within the plastic zone of the material during crack growth resistance loading. Optical microscopy of the EXCO-exposed surfaces prior to the KR loading revealed no evidence of substantial pitting and/or intergranular corrosion. Also, secondary intergranular cracking was not observed in the sodium chloride-exposed specimens, with a ductile (dimple-facet) type fracture surface directly adjacent to the exposed surface (refer to Figure 10). Additionally, microscopy of the sodium chloride-exposed alloy prior to KR loading revealed localized attack in the form of isolated pitting and intergranular corrosion on the exposed surfaces. It is thought
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that the variation in the localization of corrosive attack between the two exposure environments is a result of the difference in pH – that is, the significantly lower pH of the EXCO solution results in the rapid breakdown of the protective surface oxide layer, whereas only a localized breakdown of the oxide layer occurs at the more neutral pH of the sodium chloride solution. Consequently, uniform surface reactions may ensue in the case of the EXCO solution, while only localized (pitting) reactions are possible for the sodium chloride solution. Furthermore, it is probable that the low pH of the EXCO solution provides an abundance of hydrogen at the exposed surface, which may be absorbed via the grain boundaries of the alloy, thus embrittling The Journal of the Southern African Institute of Mining and Metallurgy
The crack growth resistance behaviour of aluminium alloy 2024-T3 at slow strain rates
Figure 8—Scanning electron fractograph showing the fracture surface appearance near the exposed surface (E) of the Exco-exposed surface with a transition from intergranular cracking (D) to ductile/dimpled shear fracture (C)
Figure 10—Scanning electron fractograph showing the fracture surface appearance near the exposed surface (E) of the sodium chloride-exposed specimen with dimpled shear fracture (C) directly adjacent to the exposed surface near corrosive attack (F)
(i)
Figure 11—MicroCT scans of the notch/crack configuration in Exco-exposed aluminium alloy 2024-T3 after crack growth resistance loading until a crack mouth opening displacement of 3 mm, with (i) showing a threedimensional MicroCT map of the notch/crack configuration and (ii) a cross-sectional view on the plane shown on the right of (i). Note that the secondary crack depth in (ii) may be correlated to the depth of the apparently intergranular fracture surface in Figure 8.
(ii) Figure 9—Higher magnification scanning electron micrographs of the fracture surface in Figure 8, with (i) the fracture surface labelled C (ductile, dimpled fracture facets), and (ii) the fracture surface labelled D (intergranular fracture) The Journal of the Southern African Institute of Mining and Metallurgy
the grain boundaries though one of the hydrogen embrittlement mechanisms mentioned in the Introduction. Further research is required to clarify if the intergranular fracture surface near the VOLUME 121
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The crack growth resistance behaviour of aluminium alloy 2024-T3 at slow strain rates EXCO-exposed surface is in fact due to hydrogen embrittlement and, if so, which of the hydrogen embrittlement mechanisms induced the embrittlement; i.e. the HEDE, HELP, AIDE, or hydride formation mechanism. Other possible causes could be microcorrosion at the grain boundaries or galvanic-type corrosion at the interfaces of the grain boundary intermetallic precipitates and the aluminium matrix. With regard to the sodium chloride-exposed material, it is possible that the degradation in the crack growth response is adversely influenced by the localized stress concentrations that accompany the isolated pitting and areas of intergranular corrosion. The notch effects that arise from the presence of these surface defects would amplify the triaxial stress ahead of the crack front, which would aid in the extension of the crack. Nevertheless, the fact that the pits and grain boundary cracks were isolated and were not observed at the precrack tip highlights the probable contribution of hydrogen embrittlement.
Conclusions
of metals and alloys by alternate immersion in neutral 3.5% sodium chloride solution. West Conshohocken, PA. ASTM International. 2018. ASTM G34-01(2018). Standard test method for exfoliation corrosion susceptibility in 2XXX and 7XXX series aluminium alloys (EXCO Test). West Conshohocken, PA. ASTM International. 2019. ASTM E561-19. Standard test method for KR curve determination. West Conshohocken, PA. Azofeifa, D., Clark, N., Amador, A., and Sáenz, A. 1997. Determination of hydrogen absorption in Pd coated Al thin films. Thin Solid Films, vol. 300. pp. 295–298. Birnbaum, H. 1990. Mechanisms of hydrogen related fracture of metals. Hydrogen Effects on Materials Behaviour. Moody, N. and Thompson, A. (eds). TMS, Warrendale, PA. pp. 639–658. Birnbaum, H., Robertson, I., Sofronis, P., and Teter, D. 1997. Mechanisms of hydrogen related fracture – A review. Corrosion-Deformation Interactions. Magnin, T. (ed.). Instutute of Materials, London. pp. 128–195. Charitidou, E., Papapolymerou, G., Haidemenopoulos, G., Hasiotis, N., and Bontozoglou, V.
Exposing aluminium alloy 2024-T3 sheet to a standard EXCO solution for a short duration (2 hours) resulted in no general corrosion. However, a number of embrittling effects were identified. During KR testing at slow strain rates, the KC_e values were reduced by 11.2 ± 0.1% compared to the equivalent value prior to exposure. After interruption of the KR testing at a crack mouth opening displacement of 3 mm, many secondary cracks parallel to the main crack front and on the exposed surface - were observed within the plastic zone. The fracture surface in this region was intergranular, while the fracture surface of the main crack front was ductile with a dimpled nature. Exposure of the same alloy to a NaCl solution for 24 hours also led to a reduction of the KC_e values during KR testing at slow strain rates. In this case, the reduction was found to be 13.0 ± 0.5%, and isolated intergranular etching and localized pitting were observed on the exposed surface.
1999. Characterization of trapped hydrogen in exfoliation corroded aluminum alloy 2024. Scripta Materialia, vol. 41. pp. 1327–1332. Dovletoglou, E., Skarvelis, P., Stergiou, V., and Alexopoulos, N. 2018. Effect of corrosion exposure on the mechanical performance of 2024 aluminum alloy electron beam welded joints. Procedia Structural Integrity, vol. 10. pp. 73–78. Kamoutsi, H., Haidemenopoulos, G., Bontozoglou, V., and Pantelakis, S. 2006. Corrosion-induced hydrogen embrittlement in aluminum alloy 2024. Corrosion Science, vol. 48. pp. 1209–1224. Lynch, S. 2011. Hydrogen embrittlement (HE) phenomena and mechanisms. Stress Corrosion Cracking. Raja, V. and Shoji, T. (eds). Woodhead Publishing, Sawston, Cambridge, UK. pp. 90–130. Milne, I., Ritchie, R., and Karihaloo, B. 2003. Hydrogen assisted damage mechanisms. Comprehensive Structural Integrity. Elsevier. pp. 69–74. Pantelakis, G., Daglaras, P., and Apostolopoulos, C. 2000. Tensile and energy density properties of 2024, 6013, 8090 and 2091 aircraft aluminum alloy after
Acknowledgements
corrosion exposure. Theoretical and Applied Fracture Mechanics, vol. 33.
The assistance of Nikos Alexopoulos and his team at the University of the Aegean in providing the blanks for the KR testing and for fruitful discussions is gratefully acknowledged. The MicroCT analyses were sperformed by Anton du Plessis of the 3D Innovation Research Group at Stellenbosch University. The assistance of Chris McDuling and Steven Masete from the Mechanical Testing Laboratory of the CSIR in performing the fatigue precracking of the KR samples is gratefully acknowledged. The work was financially supported by the Light Metals Development Network of the Department of Science and Innovation.
pp. 117–134. Petroyiannis, P., Kermanidis, A., Papanikos, P., and Pantelakis, S. 2004. Corrosioninduced hydrogen embrittlement of 2024 and 6013 aluminum alloys. Theoretical and Applied Fracture Mechanics, vol. 41. pp. 173–183. Ratke L. and Voorhees, P. 2002. Nucleation, growth and coarsening. Growth and Coarsening. Engineering Materials. Springer, Berlin, Heidelberg. https://doi. org/10.1007/978-3-662-04884-9_10 Reynolds, A. 1996. Comparison of R-Curve methodologies for ranking the toughness of aluminium alloys. Journal of Testing and Evaluation, vol. 24. pp. 406–410. Sharp, P., Cole, G, Clark, G., and Russo, G. 1998. The influence of corrosion
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u
The Journal of the Southern African Institute of Mining and Metallurgy
The use of 4-methylbenzenesulfonate ionic liquid derivatives as environmentally friendly corrosion inhibitors for mild steel in hydrochloric acid Affiliation: 1 Department of Chemistry, School of Mathematical and Natural Sciences, University of Venda, Private Bag X5050, Thohoyandou 0950, South Africa. Correspondence to: L.C. Murulana
Email:
Lutendo.murulana@univen.ac.za
Dates:
Received: 30 Aug. 2020 Revised: 13 Apr. 2021 Accepted: 13 Apr. 2021 Published: April 2021
How to cite:
Nesane, T., Mnyakeni-Moleele, S.S., and Murulana, L.C. 2021 The use of 4-methylbenzenesulfonate ionic liquid derivatives as environmentally friendly corrosion inhibitors for mild steel in hydrochloric acid. Journal of the Southern African Institute of Mining and Metallurgy, vol. 121, no. 4, pp. 159–168.
T. Nesane1, S.S. Mnyakeni-Moleele1, and L.C. Murulana1
Synopsis The effectiveness of two synthesized ionic liquids, 1-(benzyloxy)-1-oxopropan-2-aminium 4-methylbenzenesulfonate (1-BOPAMS) and 4-(benzyloxy)-4-oxobutan-1-aminium 4-methylbenzenesulfonate (4-BOBAMS), were evaluated for mild steel corrosion inhibition in 1.0 M hydro-chloric acid solution, using electrochemical impedance spectroscopy (EIS), potentiodynamic polarization (PDP), and gravimetric techniques. Organic moieties responsible for the adsorption process on mild steel surface were investigated using Fourier transform infrared spectroscopy (FTIR). Gravimetric analysis revealed that the inhibition efficiency of 1-BOPAMS and 4-BOBAMS increased with concentration, with maximum inhibition values of 90.32% and 97.91%, respectively, at the highest concentration of the inhibitors. Gibbs free energy (ΔG°ads) values indicated a strong interaction between the mild steel surface and the molecules of the ionic liquids, and that the adsorption process was spontaneous. These values also show that the inhibitive nature of ionic liquids against mild steel corrosion is caused by a mixedtype of adsorption film formed on the steel surface. The Langmuir adsorption isotherm was used to describe the adsorption of ionic liquid molecules onto the mild steel surface. Polarization curves showed that 1-BOPAMS and 4-BOBAMS have a similar effect on both the anodic and cathodic half-reactions, indicating that they prevent the dissolution of mild steel through both physical and chemical process. Nyquist plots were defined by incomplete semicircle capacitive loops, showing that the charge transfer mechanism controls the corrosion of mild steel in acidic solution. Keywords corrosion inhibition, ionic liquids, mild steel, adsorption isotherm.
DOI ID: http://dx.doi.org/10.17159/24119717/1343/2021 ORCID T. Nesane https://orchid.org/0000-00032066-1124 S.S. Mnyakeni-Moleele https://orchid.org/0000-00027124-9124 L.C. Murulana https://orchid.org/0000-00015889-3794
Introduction An inhibitor that is both effective and environmentally friendly, as well as inexpensive and which produces the desired effect at low concentrations, is preferable for use against metal corrosion. Some studies have been devoted to the subject of corrosion inhibitors (Musa, Jalgham, and Mohamad 2012; Nataraja, Venkatesha, and Tandon, 2012). Corrosion inhibitors protect against corrosion by forming protective films on the metal surface, diminishing any possible contact of the metal surface with the corrosive environment. For inhibitors to be able to protect the metal from corrosion, they must first reach the metal surface and electrochemically react with, or be adsorbed onto, the metal surface (Cisse et al., 2011; Raja et al., 2016). As such, they must have centres or functional groups in their molecules with a high electron density from which they can donate electrons to the metal surface, resulting in the coordination of the inhibitor to the metal surface (Chaubey, Quraishi, and Ebenso, 2015; Raja et al., 2016). Mild steel (MS) has many applications in industries, but readily undergoes rusting (oxidation) when exposed to corrosive environments (Kim, Kim, and Moon, 2011). Ionic liquids (ILs) are preferred as corrosion inhibitors due to their chemical structure, which contains both organic and inorganic functional groups that act as centres of adsorption onto metal surfaces (Lozano et al., 2014). Zheng et al. (2014) studied the use of ILs as corrosion inhibitors for MS in sulphuric acid, and found that they acted as an effective mixed-type inhibitor with a predominantly cathodic nature. Although the use of some ILs as corrosion inhibitors has been reported, owing to their physicochemical
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The use of 4-methylbenzenesulfonate ionic liquid derivatives properties further studies of other ILs are warranted due to the potential for discovering excellent inhibitors for metals in different corrosive environments. As such, the prime aim of the present study was to synthesize, and investigate the effectiveness of, two ILs, 4-(benzyloxy)-4-oxobutan-1-aminium 4-methylbenzenesulfonate (4-BOBAMS) and 1-(benzyloxy)-1oxopropan-2-aminium 4-methylbenzenesulfonate (1-BOPAMS), as potential inhibitors for MS corrosion in 1.0 M HCl.
Experimental procedure Materials MS with a composition of Mn 0.37 wt% , Ni 0.039 wt%, P 0.02 wt%, S 0.03 wt%, C 0.21 wt%, Fe 99.32 wt%, and Mo 0.01 wt% was used in all experimental procedures. The MS coupons had a length of 3 cm and a breadth of 2 cm and contained a small hole 2 mm in diameter for hanging on a glass rod.
Preparation of solutions 1.0 M HCl solution was prepared by diluting analytical grade (32%) HCl with double-distilled water. Stock solutions of 10.0 × 10-3 M of the two corrosion inhibitors were prepared by weighing appropriate amounts of the synthesized compounds in a 1000 ml volumetric flask. From each stock solution of the inhibitor, a series of solutions of different concentrations was prepared. All the experiments were carried out at a 1:1 ratio of 1.0 M HCl and the inhibitor of interest.
Synthesis of 1-BOPAMS and 4-BOBAMS The two IL inhibitors were synthesized by suspending alanine (6.2124 g, 69.7395 mmol) and γ-amino butanoic acid (7.0164 g, 68.0411 mmol) respectively, with 10 ml benzoyl alcohol (10.4 g, 96.1716 mmol) in a mixture of toluene ( PhMe) (200 ml) and p-toluenesulfonic acid monohydrate (PTSA) (14.5821 g, 84.6812 mmol). The mixture was refluxed for 10 hours with azeotropic removal of water. The products obtained were precipitated by the addition of diethyl ether (Et2O) (100 ml), filtered, dissolved in CH3OH (60 ml), and further precipitated by the addition of Et2O (100 ml). Figure 1 shows the reaction scheme followed for the synthesis of the two compounds. The products were obtained as white crystals, filtered, and dried. 1-BOPAMS (C17H21NO5S, mol. wt. 351.42), yield:
22.3816 g, 91% and 4-BOBAMS (C18H23NO5S, mol. wt. 365.44), yield: 23.0146 g, 93%. Spectral analysis was carried out to confirm the two ionic liquids.
Weight loss measurements The initial mass of the MS coupons was noted before they were immersed completely in 100 ml of 1.0 M HCl in the absence and presence of different concentrations of 1-BOPAMS and 4-BOBAMS and at various temperatures ranging from 30–60°C for 8 hours. This was done in a thermostatic water bath to maintain the temperature during the immersion period. After 8 hours, the specimens were removed, washed with double-distilled water, and dried. After drying, the final weights of the specimens were noted using a PW-254 Adam analytical balance at four decimal places.
Electrochemical techniques Electrochemical measurements were carried out using a Metrohm Autolab potentiostat/galvanostat (PGSTAT302N). The corrosion cell consisted of the MS as the working electrode, a platinum counter-electrode, and a saturated calomel electrode (SCE) as a reference electrode. The MS was left to corrode freely for 30 minutes to attain a stable value of Ecorr in an open circuit potential (OCP) versus SCE before any tests was started. The EIS analysis was carried out at a frequency range of 100 kHz to 0.00001 kHz under OCP conditions with an amplitude 5 mV peak to actual peak, using an AC signal at Ecorr. The PDP curves were recorded at a potential range of –250 to +250 mV (vs SCE) and a scan rate of 1 mV.s-1. The semicircle of the Nyquist plots was fitted to obtain the impedance parameters using Equation [1]: [1] where %IEEIS is the percentage inhibition efficiency from the electrochemical impedance spectroscopy data, R°ct is the charge transfer resistance in the absence of the inhibitor, and Rct is the charge transfer resistance in the presence of the inhibitor. Tafel curves (anodic and cathodic) obtained from the PDP measurements were extrapolated to obtain the electrochemical parameters. Equation [2] was used to calculate inhibition efficiencies (%IEPDP).
Figure 1—Schematic of the synthesis of 1-BOPAMS and 4-BOBAM. (1) Alanine, (2) γ-amino butanoic acid, (3) benzoyl alcohol
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The use of 4-methylbenzenesulfonate ionic liquid derivatives [2] where, %IEPDP is the percentage inhibition efficiency from the potentiodynamic polarization data, i°corr is the corrosion current density in the absence of an inhibitor, and iicorr is the corrosion current density in the presence of an inhibitor.
Results and discussion Characterization of 4-BOBAMS and 1-BOPAMS Figure 2 shows the 1H-NMR and 13C-NMR spectra of 1-BOPAMS. The 1H-NMR spectra were characterized by a broad singlet peak accounting for three protons at 8.36 ppm. The spectra also had two doublets which account for two protons each at 7.50 ppm and 7.13 ppm. The aromatic region (this is the region of the NMR
Figure 2—1H-NMR and 13C-NMR spectra of 1-BOPAMS, where (a) and (b) represent the 1H-NMR spectrum and (c) represents the 13C-NMR spectrum The Journal of the Southern African Institute of Mining and Metallurgy
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Figure 2—1H-NMR and 13C-NMR spectra of 1-BOPAMS, where (a) and (b) represent the 1H-NMR spectrum and (c) represents the 13C-NMR spectrum (continued)
spectrum where aromatic protons, which are protons attached to the benzene rings of the ILs, appear) was characterized by a multiplet accounting for five protons at about 7.4 ppm. The heteroatomic and aliphatic regions of 1-BOPAMS were characterized by a doublet accounting for three protons at 1.41 ppm, a singlet accounting for three protons at 3.41 ppm, and another singlet accounting for two protons at 5.34 ppm. The presence of a quintet accounted for single protons at around 4.18 ppm. The 13C-NMR spectra represent the accurate number of carbon atoms present in the chemical structures of the compounds used at the appropriate chemical shift values. Two methyl carbon peaks were observed at around 16.16 ppm (H3C(3)Aln) and 21.25 ppm (H3C(3)PTSA). A methine carbon peak was observed at approximately 31.16 ppm (HC(2)Aln) and a methylene carbon peak at around 67.49 ppm (H2CBn). The aromatic protons of the benzene rings were found at around 125–128 ppm (CAr), and four quintenary peaks at about 135, 138, 145, and 170 ppm (CO2Bn). 4-BOBAMS was characterized in a similar way as 1-BOPAMS and both the 1H-NMR spectra and 13C-NMR are comparable to those of 1-BOPAMS.
Electrochemical impedance spectroscopy (EIS) An electrode equivalent circuit can be fitted to interpret experimental data collected from EIS, where the individual elements of the circuit correspond to the electrochemical properties of the tested system (Macdonald, 1992); in this case, 1.0 M HCl in the absence and presence of 4-BOBAMS and 1-BOPAMS corrosion inhibitors. Figures 3a and 3b shows the blank fitted Nyquist plot and the three-element equivalent circuit model used to measure the impedance of the two ILs. The circuit included the solution resistance (Rs) represented by (R1), charge transfer resistance (Rct), represented by (R2), and (Q1) represents the constant phase element (CPE) (Yo).
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Figure 4 shows that the Nyquist plots obtained consisted of capacitive loops in the absence and presence of different concentrations of the inhibitors at 30°C. The plots are represented by incomplete semicircle capacitive loops. This imperfection indicates that the charge transfer process governs the corrosion of MS in acidic solutions (Jacob and Parameswaran, 2010) and is possibly due to factors such as the frequency dispersion, distribution of surface active sites, inhomogeneity of the metal surface, grain boundaries, surface roughness, and impurities. In such cases, CPE is usually introduced in the circuit to model a complex impendence system to get a more accurate fit (Ehsani et al., 2014a, 2014b). CPE is characterized by a fixed phase shift angle and its impedance is defined by Equation [3]: [3] where, ZCPE is the constant phase element impedance, Yo (CPE constant) and n (CPE exponent) are parameters related to phase shift angle, ω is the angular frequency (ω = 2πf, where f is the AC frequency), and j is the imaginary unit (Yadav, Quraishi, and Maiti, 2012). The significant change in the impedance response in the presence of the inhibitors can be attributed to the increase in the incomplete semicircle capacitive loop diameter with the increase in the inhibitor concentration, which specifies the increasing coverage of the MS surface. The electrochemical impedance parameters derived from Nyquist plots and the %IEEIS values are listed in Table I. The Rct values are directly proportional to the corrosion inhibition efficiency and provide details on the magnitude of the electron transfer over the surface of the MS (Jacob and Parameswaran, 2010). It can be observed from Table I that by raising the inhibitor concentration, the Rct values also increased while the CPE The Journal of the Southern African Institute of Mining and Metallurgy
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Figure 3 – Stimulated circuit fit (a) and equivalent circuit (b) used to fit the impendence data for MS corrosion in 1.0 M HCl at 30°C
Table I
E lectrochemical impedance spectroscopy (EIS) parameters including the inhibitor concentration (Cinh), resistance of charge transfer (Rct), constant phase element (Yo), solution resistance (Rs), CPE exponent (n), and the surface coverage (θ) for MS corrosion in 1.0 M HCl in absence and presence of different concentrations of 1-BOPAMS and 4-BOBAMS at 30°C Inhibitor
Cinh (M)
Rct (Ω)
Yo (F.s(a-1))
Rs (Ω)
n
θ
IEEIS (%)
Blank
– 1.0 x10–3 2.0 x10–3 3.0 x10–33
4.37 9.98 16.06 17.75
0.26e-3 0.25e-3 0.181e-3 0.26e-3
1.867 2.314 2.665 6.478
0.8827 0.8787 0.8551 0.8729
– 0.5627 0.7282 0.7541
– 56.27 72.82 75.41
1-BOPAMS
4.0 x10–3 5.0 x10–3 1.0 x10–3 2.0 x10–3 3.0 x10–3
24.11 57.50 17.57 17.75 19.46
0.15e-3 0.14e-3 0.20e-3 0.26e-3 0.16e-3
2.360 2.827 2.391 6.478 3.578
0.8795 0.8958 0.8833 0.8729 0.8938
0.8190 0.9241 0.7541 0.7541 0.7757
81.90 92.41 75.41 75.41 77.57
4-BOBAMS
4.0 x10–3 5.0 x10–3
58.05 94.62
0.22e-3 0.22e-3
3.950 2.885
0.8352 0.8571
0.9248 0.9539
92.48 95.39
Figure 4 – Nyquist plots for MS in 1.0 M HCl in the absence and presence of different concentrations of 1-BOPAMS and 4-BOBAMS at 30°C
decreased. This is because the addition of inhibitor increased the surface coverage on the MS by the inhibitor molecules, and due to the formation of a protective film which resulted in a decrease of the electron transfer between the metal surface and the The Journal of the Southern African Institute of Mining and Metallurgy
corrosive medium (Chetouani, 2002). The CPE constant for both inhibitors at all studied concentrations decreased compared to the blank. The decrease in Yo after the addition of the inhibitors may be due to either an increase in the thickness of the double VOLUME 121
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The use of 4-methylbenzenesulfonate ionic liquid derivatives layer (this is a layer that is formed on the surface of the metal when it is in contact with the corrosive solution in the presence of the inhibitors) or to the desorption of water molecules from MS followed by the adsorption of the inhibitor onto the metal surface. The CPE indicates a reduction in the local dielectric constant or an enhancement in the thickness of the electric double layer.
Potentiodynamic polarization (PDP) PDP measurements were carried out to understand the role of inhibitors in biasing anodic and cathodic reactions of MS. Before running any tests, the system was allowed to reach a constant potential at an open circuit. Anodic and cathodic Tafel curves for MS corrosion in 1.0 M HCl solution were obtained in the absence and presence of various concentrations of 4-BOBAMS and 1-BOPAMS at 30°C (Figure 5). Potentiodynamic parameters such as corrosion potential (Ecorr), corrosion current density (icorr), anodic (βa) and cathodic (βc) Tafel slopes, and linear polarization resistance were extrapolated from these curves, and the results are shown in Table II. The Tafel curves and Table II show that the icorr decreased markedly upon introduction of the two inhibitors both in the anodic and cathodic regions, indicating that the inhibitors are adsorbed onto the MS surface and inhibit the corrosion process. An inhibitor can be classified as anodic or cathodic type if the change in the Ecorr values is greater than ±85 mV (Mashuga et al., 2015). In the present study, the magnitude of the shifted Ecorr values in the presence of ILs compounds was towards
more negative values and the shift was less than 85 mV at all inhibitor concentrations with respect to the blank solution, which implies that these are mixed-type inhibitors that control both the cathodic and anodic reactions of the MS. Nesane, MnyakeniMoleele, and Murulana (2020) observed a similar trend in the shift of Ecorr values caused by the same inhibitors against aluminium corrosion. On the other hand, the addition of different concentrations of these inhibitors in 1.0 M HCl solution altered the Tafel slopes (βa and βc) to approximately the same extent, supporting the concept that both the retardation of the anodic metal dissolution and cathodic hydrogen reduction were affected (Hegazy, Ahmed, and El-Tabei, 2011). The variation of βa and βc in the presence of the inhibitors compared to the blank can be attributed to the change in the kinetics of hydrogen evolution as a result of the diffusion or barrier effect (Tourabi et al., 2013). The cathodic polarization curves are almost parallel to each other, indicating that the hydrogen evolution reaction is under activation control (Thanapackiam et al., 2016). The change in βa values with the addition of the inhibitors suggests that the inhibition process could be attributed to the formation of an adsorption film on the metal surface, thus impeding the corrosion of MS by blocking the active sites of the metal without affecting the anodic reaction mechanism (Abdel-Rehim, Khaled, and Al-Mobarak, 2011). The addition of the ILs surfactants to the corrosive solution resulted in an increase in the Rp values, which increased to a maximum at the highest concentration of the surfactants used in
Figure 5—Tafel plots for MS in 1.0 M HCl in the absence and presence of different concentrations of 1-BOPAMS and 4-BOBAMS inhibitor compounds (shown by different colours)
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P otentiodynamic polarization (PDP) parameters including the inhibitor concentration (Cinh), corrosion potential (Ecorr), corrosion current density (icorr) and anodic and cathodic Tafel slopes (βa and βc) for MS corrosion in 1.0 M HCl in both the presence and absence of different concentrations of 1-BOPAMS and 4-BOBAMS at 30°C
Inhibitor
Cinh (M)
Ecorr (mV)
icorr (mAcm-2)
βa (mVdec-1)
βc (mVdec-1)
Rp (10-1) (Ω cm2)
Blank
– –469 1.0 x10–3 –477 2.0 x10–3 –483
0.00212 0.00067 0.00023
20 8 5
26 10 6
22.97 28.33 54.60
IEPDP (%)
– – 68.55 78.63 89.40 83.43
IEWL (%)
1-BOPAMS
3.0 x10–3 4.0 x10–3 5.0 x10–3 1.0 x10–3
–447 –478 –468 –481
0.00030 0.00023 0.00017 0.00045
9 8 7 7
13 11 10 11
79.69 87.46 100.34 41.04
86.33 84.77 89.11 88.75 91.92 90.32 78.67 92.73
4-BOBAMS
2.0 x10–3 3.0 x10–3 4.0 x10–3 5.0 x10–3
–486 –480 –475 –478
0.00045 0.00033 0.00026 0.00024
6 7 8 7
8 10 12 10
33.40 52.38 82.16 76.82
78.71 94.51 84.28 97.65 87.75 97.70 88.67 97.91
Table III
W eight loss measurements of MS in 1.0 M HCl containing various concentrations of 1-BOPAMS and 4-BOBAMS at different temperatures (where Cinh stands for inhibitor concentration, CR stands for corrosion rate, and IE stands for inhibition efficiency) Inhibitor
Blank
30°C
Cinh (M)
CR (g.cm-2.h-1)
40°C
IE (%)
CR (g.cm-2.h-1)
50°C
IE (%)
CR (g.cm-2 h-1)
60°C
IE (%)
CR (g.cm-2.h-1)
IE (%)
–
0.00717
–
0.01320
–
0.02437
–
0.04099
–
1-BOPAMS
1.0×10-3 2.0×10-3 3.0×10-3 4.0×10-3 5.0×10-3
0.00153 0.00189 0.00109 0.00081 0.00069
78.63 83.43 84.77 88.75 90.32
0.00432 0.00404 0.00399 0.00383 0.00340
67.27 69.37 69.80 71.01 74.23
0.00909 0.00860 0.00848 0.00733 0.00715
62.72 64.73 65.19 69.92 70.68
0.01797 0.01725 0.01692 0.01650 0.01569
56.17 57.93 58.73 59.75 61.73
4-BOBAMS
1.0×10-3 2.0×10-3 3.0×10-3 4.0×10-3 5.0×10-3
0.00052 0.00039 0.00017 0.00017 0.00015
92.73 94.51 97.65 97.70 97.91
0.00296 0.00283 0.00261 0.00205 0.00180
77.59 78.54 80.24 84.49 86.40
0.00990 0.00931 0.00898 0.00882 0.00833
59.38 61.81 63.14 63.80 65.81
0.01880 0.01899 0.01776 0.01686 0.01576
53.28 54.13 56.68 58.88 61.56
this study, indicating an effective inhibition by the compounds. The increase in the Rp values with the increase in inhibitor concentration implies that further polarization of MS was opposed by the formation of an adsorption film formed by the inhibitor molecules present in the solution at the metal/solution interface. The %IEPDP values obtained increased with the increased concentration of both inhibitors.
Effect of temperature and inhibitor concentration The corrosion rate (CR) and %IE in the absence and presence of various concentrations of 4-BOBAMS and 1-BOPAMS in 1.0 M HCl solution and at different temperatures (30–60°C) are presented in Table III. The addition of 4-BOBAMS and 1-BOPAMS markedly decreased the corrosion rate of MS and the decrease was proportional to the increase in the concentration of the inhibitors. In the uninhibited system, the weight loss was higher and decreased at a faster rate as the temperature of the solution was increased from 30°C to 60°C. The acceleration in the weight loss can be attributed to the increase in collisions between the HCl molecules and the MS surface. However, when an inhibitor was introduced into the solution, the weight loss decreased significantly; this was due to the formation of a surface film by the adsorption of the inhibitor molecules, which acted as a barrier between the aggressive HCl molecules and MS surface. Even in the presence of the inhibitors in solution, the %IE decreased as the temperature was increased, possibly due to the The Journal of the Southern African Institute of Mining and Metallurgy
desorption of the inhibitors on the MS surface. For instance, the %IE of 1-BOPAMS at 30°C is 90.32%, decreasing to 61.73% at 60°C for the highest concentration of 5.0 ×10-3 M. This behaviour was observed for both inhibitors. The protective properties of the inhibitors are possibly a result of the interaction between π-electrons and heteroatoms (π-electrons are electrons that are involved in a π-bond between two atoms (this can be in form of double or triple bonds), and heteroatoms are any atoms other than carbon or hydrogen, such as sulphur, oxygen, and nitrogen, that are in the two chemical structures of the compounds utilized in this study) with the positively charged MS surface (Mu et al., 2006).
Adsorption considerations The data obtained from the weight loss analysis was fitted to several isotherms. The Langmuir adsorption isotherm gave the best description for the synthesized compounds on MS in 1.0 M HCl solution. The Langmuir adsorption isotherm is calculated from Equation [4]. [4]
where Cinh is the concentration of the inhibitor, Kads is the equilibrium constant of adsorption, and θ is the surface coverage. VOLUME 121
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The use of 4-methylbenzenesulfonate ionic liquid derivatives The plots of Cinh/θ versus Cinh gave a straight line (Figure 6) with a regression coefficient (R2) around unity, confirming that the adsorption of 4-BOBAMS and 1-BOPAMS on a MS surface in 1.0 M HCl obeyed the Langmuir adsorption isotherm. From these figures, the equilibrium constant of adsorption (Kads) was obtained from the slopes and enabled the calculation of the free Gibbs energy using Equation [5]. [5] where, the value 55.5 is the molar concentration of water in the solution, ΔG°ads is the standard Gibbs free energy of adsorption, R is the universal gas constant, T is the absolute temperature (all in metric units), and Kads is the equilibrium constant of adsorption. The values of ΔG°ads around –0 kJ.mol-1 and lower are consistent with the electrostatic interaction between charged molecules and the charged surface of the substrate (physisorption), which is a process that involves weak Van der Waals interaction rather than chemical bonding (Musa, Jalgham, and Mohamad, 2012). Conversely, those that are around –40 kJ.mol-1 and above correspond to chemical adsorption (chemisorption), which occurs through the sharing or transfer of an electron from the adsorbate molecules on the surface of MS resulting in a coordinate type of bond. The ΔG°ads values (Table IV) were negative, indicating the spontaneity of the adsorption process and that there was a strong interaction between the inhibitor molecules and MS surface (Hegazy, Ahmed, and ElTabei, 2011). The ΔG°ads values that are around –20 kJ.mol-1, and also those around –40 kJ.mol-1, implied that mixed-type
adsorption took place, i.e., a mixture of both physisorption and chemisorption. The trends of these values elaborate on the nature of the adsorption that occurs between the MS surface and the corrosion inhibitors. Thiraviyam and Kannan (2013) found that increased ΔG°ads with increased temperature indicates exothermic adsorption, whereas decreasing ΔG°ads with increased temperature indicates endothermic adsorption. In the present study, the ΔG°ads values increased as the temperature increased, indicating that the adsorption mechanism between the MS surface and the two inhibitors was exothermic, which signifies either chemisorption or physisorption mechanisms.
Adsorption film analysis Figure 7 reveals that some of the functional groups observed in the pure compounds disappeared, and these functional groups can be said to be responsible for the complex formation with the MS surface, preventing the dissolution process. The FTIR spectra for the adsorption film formed show that the intensity of the peaks decreased, indicating that a coordinate bond was created through the functional groups of these peaks, with the Fe2+, forming the Fe2+-inhibitor complex on the surface of the MS, preventing the dissolution process. The FTIR spectra for the two pure ILs, 4-BOBAMS and 1-BOPAMS, were similar and showed an ammonium salt (+NH3) which gave a strong, broad absorption band at around 3066–3095 cm-1. The intensity of these peaks decreased and shifted to about 3326 cm-1. The peaks located around 630–660 cm-1 represent the Fe-O stretching bond frequency, indicating the modification of the MS surface by the
Figure 6—Langmuir adsorption isotherm plots for the adsorption of different concentrations of 1-BOPAMS and 4-BOBAMS on the surface of MS in 1.0 M HCl at different temperatures
Table IV
Thermodynamic and adsorption parameters for MS in 1.0 M HCl at various temperatures for 1-BOPAMS and 4-BOBAMS
Inhibitor
T (K)
1-BOPAMS
303 313 323 333
0.9985 0.9971 0.9962 0.9985
1.0638 1.3249 1.3590 1.5881
3864.29 4296.82 2949.16 3806.91
–30.9266 –32.2234 –32.2422 –33.9473
4-BOBAMS
303 313 323 333
0.9999 0.9976 0.9992 0.9966
1.0035 1.1184 1.4861 1.5589
12025.01 3790.46 4009.78 2086.51
–33.7866 –31.8971 –33.0673 –32.2823
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Figure 7—FTIR spectra for the pure compound and adsorption films formed on the MS in 1.0 M HCl by 1-BOPAMS and 4-BOBAMS corrosion inhibitors
formation of a Fe-inhibitor complex through the oxygen atoms present in the structures of the compounds. The peaks around 3326.22 cm-1 and 1614.52 cm-1 for MS represent the hydrated iron oxide molecule of (Fe2O3.xH2O) and the presence of CO group in the form of HCO3- ion participating during corrosion, which is in agreement with findings by Renita (2015).
Conclusions The 1-BOPAMS and 4-BOBAMS were successfully synthesized and evaluated as corrosion inhibitors of mild steel in HCl acidic medium. The inhibitors are adsorbed onto the mild steel surface by a combination of both physisorption and chemisorption (i.e., mixed-type adsorption); however, chemisorption was the dominant adsorption process. Nyquist curves are represented by an imperfect semicircle loop, indicating effective inhibition of mild steel corrosion. The charge transfer resistance increased with inhibitor concentration, leading to a decrease of the corrosion rate and higher percentage inhibition efficiency. The increase of the RP values with increasing concentration of the inhibitors indicates that the inhibition of corrosion was primarily due to protective film formation.
Conflicts of Interest The authors declare no conflicts of interest regarding the publication of this paper.
Acknowledgements The authors would like to thank Sasol University Collaboration Grant, NRF and Sasol Foundation for funding this project.
References
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effect of polypyrrole films. Anti-Corrosion Methods and Materials, vol. 61, no. 3. pp. 146–152. Hegazy, M.A., Ahmed, H.M., and El-Tabei, A.S. 2011. Investigation of the inhibitive effect of p-substituted 4-(N, N, N-dimethyldodecylammonium bromide) benzylidene-benzene-2-yl-amine on corrosion of carbon steel pipelines in acidic medium. Corrosion Science, vol. 53, no. 2. pp. 671–678. Jacob, K.S. and Parameswaran, G. 2010. Corrosion inhibition of mild steel in hydrochloric acid solution by Schiff base furoin thiosemicarbazone. Corrosion Science, vol. 52, no. 1. pp. 224–228. Kim, S., Kim, J., and Moon, I. 2011. Parameter-based model for the forecasting of pipe corrosion in refinery plants. Industrial and Engineering Chemistry Research, vol. 50, no. 22. pp. 12626–12629. Lozano, I., Mazario, E., Olivares-Xometl, C.O., Likhanova, N.V., and Herrasti, P. 2014. Corrosion behaviour of API 5LX52 steel in HCl and H2SO4 media in the presence of 1, 3-dibencilimidazolio acetate and 1, 3-dibencilimidazolio dodecanoate ionic liquids as inhibitors. Materials Chemistry and Physics, vol. 147.s pp. 191–197. Macdonald, J.R. 1992. Impedance spectroscopy. Annals of Biomedical Engineering, vol. 20, no. 3. pp. 289–305. Mashuga, M.E., Olasunkanmi, L.O., Adekunle, A.S., Yesudass, S., Kabanda, M.M., and Ebenso, E.E. 2015. Adsorption, thermodynamic and quantum chemical studies of 1-hexyl-3-methylimidazolium based ionic liquids as corrosion inhibitors for mild steel in HCl. Materials, vol. 8, no. 6. pp. 3607–3632. Mu, G., Li, X., Qu, Q., and Zhou, J. 2006. Molybdate and tungstate as corrosion inhibitors for cold rolling steel in hydrochloric acid solution. Corrosion Science, vol. 48, no. 2. pp. 445–459. Musa, A.Y., Jalgham, R.T., and Mohamad, A.B. 2012. Molecular dynamic and quantum chemical calculations for phthalazine derivatives as corrosion inhibitors of mild steel in 1 M HCl. Corrosion Science, vol. 56. pp.176–183. Nataraja, S.E., Venkatesha, T.V., and Tandon, H.C. 2012. Computational and experimental evaluation of the acid corrosion inhibition of steel by tacrine. Corrosion Science, vol. 60. pp. 214–223. Nesane, T., Mnyakeni-Moleele, S.S., and Murulana, L.C. 2020. Exploration of synthesized quaternary ammonium ionic liquids as unharmful anti-corrosives for aluminium utilizing hydrochloric acid medium. Heliyon, vol. 6, no. 6. pp. e04113–e04113. Raja, P.B., Ismail, M., Ghoreishiamiri, S., Mirza, J., Ismail, M.C., Kakooei, S., and Rahim, A.A. 2016. Reviews on corrosion inhibitors: a short view. Chemical Engineering Communications, vol. 203, no. 9. pp. 1145–1156. Renita, D., Sanish, T., Dwivedi, P., and Amit, C. 2015. Green approach to corrosion inhibition by Emblica officinalis (NA-7) leaves extract. International Journal of Nano Corrosion Science and Engineering, vol. 2, no. 3. 29–45. Thanapackiam, P., Rameshkumar, S., Subramanian, S.S., and Mallaiya, K. 2016. Electrochemical evaluation of inhibition efficiency of ciprofloxacin on the corrosion of copper in acid media. Materials Chemistry and Physics, vol. 174. pp. 129–137. Thiraviyam, P. and Kannan, K. 2013. Inhibition of aminocyclohexane derivative on mild steel corrosion in 1 N HCl. Arabian Journal for Science and Engineering, vol. 38, no. 7. pp. 1757–1767. Tourabi, M., Nohair, K., Traisnel, M., Jama, C., and Bentiss, F. 2013. Electrochemical and XPS studies of the corrosion inhibition of carbon steel in hydrochloric acid pickling solutions by 3, 5-bis (2-thienylmethyl)-4-amino-1, 2, 4-triazole. Corrosion Science, vol. 75. pp. 123–133. Yadav, D.K., Quraishi, M.A., and Maiti, B. 2012. Inhibition effect of some benzylidenes on mild steel in 1 M HCl: An experimental and theoretical correlation. Corrosion Science, vol. 55. pp. 254–266. Zheng, X., Zhang, S., Gong, M., and Li, W. 2014. Experimental and theoretical study on the corrosion inhibition of mild steel by 1-octyl-3-methylimidazolium L-prolinate in sulfuric acid solution. Industrial and Engineering Chemistry Research, vol. 53, no. 42. pp. 16349–16358. u VOLUME 121
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Deformation and fracture behaviour of the g-TiAl based intermetallic alloys M.N. Mathabathe1,2, A.S. Bolokang1, G. Govender1, C.W. Siyasiya2, and R.J. Mostert2
Affiliation: 1 Council for Scientific and Industrial Research, Manufacturing Cluster, Advanced Materials Engineering, Meiring Naude Road, P O Box, 395, Pretoria, South Africa. 2 Department of Materials Science and Metallurgical Engineering, Faculty of Engineering, Built Environment and Information Technology, University of Pretoria, South Africa. Correspondence to: M.N. Mathabathe
Synopsis The b-solidifying g-TiAl intermetallic alloys of nominal composition Ti-48Al (binary alloy), Ti-48Al-2Nb (ternary alloy), Ti-48Al-2Nb-0.7Cr (quaternary alloy), and Ti-48Al-2Nb-0.7Cr-0.3Si (quinary alloy) (in at.%) were developed. The materials produced were tensile tested at room temperature. The as-cast microstructures and fracture surfaces of the tensile tested specimens were examined using conventional metallographic methods. Microstructural examination indicated that the alloys were comprised of lamellar structures (a2+g) embedded in columnar dendritic cores in the as-cast condition. However, the quinary alloy contained a Ti5Si3 second phase. The alloys exhibited no detectable ductility during tensile deformation, indicating the brittleness of all the materials. The fracture surfaces revealed that the alloys failed by translamellar fracture with correspondingly few cleavage facets. Keywords g-TiAl based alloys, b-solidifying, translamellar fracture.
Email:
nmathabathe@csir.co.za
Dates:
Received: 30 Aug. 2020 Revised: 26 Mar. 2021 Accepted: 1 Apr. 2021 Published: April 2021
How to cite:
Mathabathe, M.N., Bolokang, A.S., Govender, G., Siyasiya, C.W., and Mostert, R.J. 2021 Deformation and fracture behaviour of the g-TiAl based intermetallic alloys. Journal of the Southern African Institute of Mining and Metallurgy, vol. 121, no. 4, pp. 169–174. DOI ID: http://dx.doi.org/10.17159/24119717/1344/2021 ORCID M.N. Mathabathe https://orchid.org/0000-00017058-5647
Introduction The advent of high-temperature materials with potential use as structural materials, shape memory alloys, and coatings is of significance for a range of applications. Single-crystal nickel-based superalloys are the most successful gas turbine materials as a result of their good mechanical properties at elevated temperature imparted by the g′-Ni3Al (L12) phase in the fcc matrix (Liu et al., 2017). The g-TiAl intermetallic alloys are competitive materials for elevated temperature applications, mainly because of reduced density, high specific strength, and modulus (Mathabathe et al., 2018a, 2018b). The commercialization of g-TiAl intermetallic alloys depends on microstructure and texture control by thermomechanical processing techniques (Mathabathe et al., 2019b). Research interest in titanium aluminides was stimulated due to their outstanding combination of high melting temperature and low density with desirable medium/high temperature mechanical properties and notable corrosion and oxidation resistance. These properties made the a2/g-based titanium aluminides suitable for application in many fields (Grytsiv et al., 2005; Kumaran et al., 2008). These intermetallic ordered compounds are stiffer, lighter, and have better mechanical properties than conventional titanium alloys, with relatively low ambient-temperature tensile strength (Lipsitt, Shechtman, and Schafrik, 1980). Yan et al. (2009) reported that the TiAl-based alloys have a higher chemical reactivity, larger solidification shrinkage rate, and lower castability than typical Al casting alloys. This can cause misrun defects on the surface of the components, but also results in shrinkage porosity and cracking defects due to poor ductility. Therefore, further alloying of TiAl alloys made sense as this served to enhance the mechanical properties. Lipsitt, Shechtman, and Schafrik (1980) studied the stoichiometric behaviour of intermetallic compounds in TiAl alloys to ascertain the potential for enhancing the ductility-temperature relationship. The study reported on the tensile and fracture properties over a wide temperature range. The results showed a significant relationship between dislocation structures, fracture modes, and tensile properties at each temperature. In our previous study (Mathabathe et al., 2018b) the effects of alloying elements on the grain size and macrohardness were demonstrated. The study indicated that the addition of Nb and Cr played an important role in decreasing the grain size of the ternary and quaternary alloys. The binary alloy had the largest grain size of 280 µm, while alloying TiAl with 2 at.% Nb reduced the grain size in the ternary alloy to 180 µm. Elements such as Nb are known to act as β-phase stabilizers in Ti; however, in TiAl alloy Nb plays two roles – it reduces the grain size and stabilizes the β-phase. Addition of Cr improved the grain refinement further, leading to a quaternary alloy. In the quinary alloy the grain size slightly increases due to the addition of Si, implying that the ductility of the alloy was enhanced following the Hall Petch relationship (Chen et al., 2004).
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Deformation and fracture behaviour of the g-TiAl based intermetallic alloys The tensile properties of the alloys were similar to those reported by Mathabathe et al., (2018b), with changes effected by the alloying elements. The addition of Nb to binary TiAl slightly increased the tensile strength. However, the quaternary alloy showed the lowest tensile strength, due to alloying with Cr. Super-dislocation and twin nucleation are often competing mechanism in g-TiAl alloys, due to stacking fault energy (Huang and Hall, 1991). Duplex structures favour twin deformation as a result of the lower fault energy due to the presence of a2 and/or lower Al concentration with the ability to absorb and reduce the interstitial content in the g-phase (Liu et al., 2016). Therefore, in the current study we investigated the deformation behaviour of the previously studied g-TiAl-based alloys (Mathabate et al., 2018b), where the microstructure and fracture analysis after tensile testing are discussed. Moreover, the study aims to ascertain the effect of β-phase on room temperature ductility, whereby the tensile fracture is evaluated.
Experimental Pure metallic powders of Ti, Al, Nb, Cr, and Si were mixed together to give nominal compositions of Ti-48Al, Ti-48Al-2Nb, Ti-48Al-2Nb-0.7Cr, and Ti-48Al-2Nb-0.7Cr-0.7Si (in atomic percentage, at.%). These compositions were consolidated and compacted prior to vacuum arc melting. The fabrication method was adapted from Mathabathe et al. (2019a). Conventional metallographic preparation techniques were used to prepare the samples. Optical microscopy and scanning electron microscopy
(SEM), using a JEOL® JSM-6510 equipped with backscattered electron (BSE) mode and energy-dispersive X-ray spectroscopy (EDS) were used to investigate the microstructures. The EDS analysis position had a working distance (WD) of 10 mm and a take-off angle of 35° for X-ray signals was utilized, with efficient mapping analysis under secondary electron imaging (SEI) observation. The acceleration voltage employed was 20 kV (Mathabathe et al., 2018b). The morphologies of the fracture surfaces, as well as electropolished longitudinal sections of the tensile specimens, were investigated using SEM-SEI (secondary electron imaging). Room-temperature tensile tests using an MTS criterion C45 tensile machine were conducted on flat tensile specimens of dimensions 5 mm × 3 mm cross-section and 12 mm gauge length, using a constant crosshead speed of 0.25 mm/min.
Results and discussion Microstructure analysis Figure 1 shows the optical microstructures of the binary, ternary, quaternary, and quinary alloys. The microstructures in Figures 1a, 1b, and 1c are comprised of the light network primary β-phase, while the dark features are g-phase. Due to the addition of Si, the sample in Figure 1d shows the light phase as quinary alloy due to the presence of titanium silicate precipitate (Ti5Si3). The β-grains are irregular in shape, as well as the fine g-grains along the grain boundaries. The microsegregation of β-stabilizers Nb and Cr on the dendritic cores possibly formed as a result of
Figure 1—Optical images of the (a) binary (Ti-48Al), (b) ternary (Ti-48Al-2Nb), (c) quaternary (Ti-48Al-2Nb-0.7Cr), and (d) quinary (Ti-48Al-2Nb-0.7Cr-0.3Si) alloys
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Deformation and fracture behaviour of the g-TiAl based intermetallic alloys solidification, while few retained primary b-phases were observed after cooling due to enrichment in b-phase stabilizers (Wang et al., 2011). The SEM images of the binary, ternary, quaternary and quinary alloys are shown in Figure. 2. Figure 2a shows the microstructure of binary TiAl, comprised of the lamellar structure (a2+g). The corresponding chemical compositions of the lamellar structure of the binary, ternary, quaternary, and quinary alloys are presented is Table I. The lamellar colonies are evident, while the b-phase was segregated along the grain boundaries of the g-phase. Figures 2b-d confirm that the lamellar colonies appear finer in the alloys containing Nb, Cr, and Si. It is important to note that the Cr atoms occupy Ti-sites in a single g-phase alloy, resulting in a small Ti-Al bond adjustment and increased plasticity (Liu et al., 2016). However, when Cr atoms occupy the Al sites in a duplex alloy, the overall covalency of the Ti-Al bond is lowered and plastic deformation is favoured, decreasing the ultimate tensile strength in the quaternary alloy. Furthermore, Si addition drastically increased the tensile strength of the quinary alloy. The improved properties in the quinary alloy indicate that the addition of Si to the TiAl-based alloy promoted the room temperature ductility due to Ti5Si3 whisker formation during solidification (Sun and Froes, 2002). The interface
between the matrix and strip/rod-like Ti5Si3 particles retards the motion of dislocations, thus increasing the strength (Liu et al., 2016).
Fracture behaviour The SEM secondary electron (SE) mode images of the fracture surface after room temperature tensile tests are shown in Figure 3. A mixture of grain boundary and transgranular fracture for the binary alloy is shown in Figure 3a. The duplex structure (a2+g) underwent cleavage or lamellar interface failure. Figures 3a--c illustrate fracture surfaces with a slightly rough appearance and irregular arrangement. The planar regions are illustrated by the arrows. The addition of Cr increased the probability of intragranular failure in the duplex structure (Huang and Hall, 1991). The quinary alloy (Figure 3d) shows the flattest/rough fractured surface discontinuities. The twinning and glide at ambient temperature arise effortlessly in grains that are globular due to fewer deformation obstacles present. The dislocation movements are hindered by a variety of interfaces (short-range) resulting in limited overall deformation (Kabir et al., 2015). Figures 4a and 4b illustrate the fracture surfaces of the binary and quaternary alloys with the emphasis on showing the facets. These two alloys exhibited the lowest tensile strength
Figure 2—SEM/BSE images of the (a) binary (Ti-48Al), (b) ternary (Ti-48Al-2Nb), (c) quaternary (Ti-48Al-2Nb-0.7Cr), and (d) quinary (Ti-48Al-2Nb-0.7Cr-0.3Si) alloys
Table I
E DS chemical composition of the lamellar structure of the Ti-48Al, Ti-48Al-2Nb, Ti-48Al-2Nb-0.7Cr and Ti-48Al-2Nb0.7Cr-0.3Si alloys Alloy Binary Ternary Quaternary Quinary
Ti
Al
Nb
Cr
Si
57.13±1.58 48.87±0.46 49.45±0.89 49.39±1.28
42.87±0.73 49.42±1.59 49.45±0.34 49.39±1.28
– 1.70±0.81 1.94±1.19 1.72±1.30
– – 0.75±0.25 0.69±0.43
– – – 0.15±0.25
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Figure 3—Fracture surfaces after tensile testing of the (a) Ti-48Al, (b) Ti-48Al-2Nb, (c) Ti-48Al-2Nb-0.7Cr), and (d) Ti-48Al-2Nb-0.7Cr-0.3Si alloys
Figure 4—Fracture surfaces showing facets on the low tensile strength (a) Ti-48Al, (b) Ti-48Al-2Nb, (c) Ti-48Al-2Nb-0.7Cr), and (d) Ti-48Al-2Nb-0.7Cr-0.3Si alloys, and river patterns in all the alloys (c-f)
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Deformation and fracture behaviour of the g-TiAl based intermetallic alloys of the four alloys. Infrequent interlamellar fractures and few facets are evident in these alloys. The fracture, comprising more facets/irregularities, suggests lower plasticity, which may indicate the brittleness of the alloy, resulting in improved tensile strength. Furthermore, the fracture surfaces of all specimens, including the ternary and quinary alloys, display features such as river-like patterns characteristic of cleavage fracture (Wang et al., 2017) shown in Figures 4c-f. The cleavage/translamellar fracture originates at the particular crystallographic plane and proceeds at a rapid rate (Wang et al., 2017). A cleavage facet is a crystallographic plane with the lowest surface energy (Wang et al., 2017). In polycrystalline metals, the low-index crystallographic planes are typically not parallel, whereas when the cracks propagate cleavage steps are generated across the grains, while the river-like patterns are the top view of the cleavage steps (Wang et al., 2017). The ‘river’ patterns commence from the grain boundary flaws and terminate at the grain boundary of the same grain. Therefore, cracks originated inside the specimens. Additionally, secondary cracking at lamellar interfaces at room temperature is often observed in the ternary and the quinary alloy, as illustrated in Figures 5a and 5b. The cracks began at the grain boundary and extended by interlamellar propagation. The arrows indicate interlamellar secondary cracking. Secondary crack branching, which propagated along lamellar interfaces (which is the path of least resistance), indicates that the lamellar structure of the alloys exhibited robust resistance to translamellar fracture. Consequently, the ternary and quinary alloys had higher tensile strength than the binary and the quaternary alloys. Figures 6a-d show optical micrographs of the side views of the fractured alloy samples to further validate the fracture analysis results. The fracture zone near the surface was observed for the four alloys investigated. The microcracks are caused by the presence of the b-phase for binary, ternary, and quaternary alloys (Figures 6a-c), while for the quinary alloy (Figure 6d) the Ti5Si3 phase acted as a crack initiation site. These cracks were initiated along the boundaries between g-grains, at the triple junctions of colonies/grains and g-grains/lamellar colonies (Chen et al., 2004). In addition, the crack opening distance and length of the microcracks suggest that the grain boundaries between equiaxed g-grains and lamellar colonies act as potential sites for crack initiation. However, when the main crack stops at the grain boundary barrier, a new crack can start along the lamellar interface at the rear end of the grain barrier. This occurs when a
lamella within a grain is inclined at a large angle or perpendicular to the crack propagation plane, whereby the resulting crack is averted to delamination mode (Figures 6a-c) (Chen et al., 2004). The fracture behaviour indicated that the β-phase is brittle at room temperature. According to Hamza, Kanniah, and Harun (2009) this is due to the limited number of dislocations. The cleavage-like fracture surfaces of the studied g-TiAl alloys subsequent to room-temperature tensile tests show the brittle characteristics of the β-phase. On the one hand, the character of the β-phase has a significant beneficial effect on the mechanical properties at the intermediate (600–700°C) operating temperatures of low-pressure turbine blades, for example.
Conclusions The deformation and fracture behaviour of g-TiAl-based alloys were investigated and the following conclusions were drawn. ➤ Microstructural examination indicated that all the alloys are comprised of lamellar structures (a2+g) embedded in columnar dendritic cores in the as-cast condition. However, the quinary alloy exhibited the Ti5Si3 second phase. ➤ The binary and quaternary alloys showed more facets, on the fracture surfaces suggesting lower plasticity, which is an indication of the brittleness of the alloy and in turn corresponds to a lower tensile strength. ➤ Infrequent interlamellar fracture and few facets were evident in the alloys. The lamellar structure of the ternary and quinary alloys exhibited better strength against translamellar fracture, confirmed by the higher tensile strength of these alloys. ➤ The fracture behaviour, viz. the cleavage-like fracture surfaces of the studied g-TiAl alloys subsequent to room temperature tensile testing, shows the brittle characteristics of the β-phase.
Acknowledgement The authors thank the Department of Science and Technology (DST) and the CSIR for funding this work.
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Figure 5—Translamellar fracture with secondary cracking in (a) Ti-48Al-2Nb and (b) Ti-48Al-2Nb-0.7Cr-0.3Si alloys The Journal of the Southern African Institute of Mining and Metallurgy
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Figure 6—Optical micrographs taken near the fractured surface with subsequent room temperature tensile failure of (a) Ti-48Al, (b) Ti-48Al-2Nb, (c) Ti-48Al-2Nb0.7Cr), and (d) Ti-48Al-2Nb-0.7Cr-0.3Si alloys
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Evaluation of spheroidized tungsten carbide powder produced by induction plasma melting R.M. Dire1, H. Bissett2, D. Delport1, and K. Premlall1 Affiliation: 1 Tshwane University of Technology, Pretoria West, Pretoria, South Africa. 2 The South African Nuclear Energy Corporation SOC Ltd. (Necsa), North West Province, South Africa. Correspondence to: H. Bissett
Email:
hertzog.bissett@necsa.co.za
Dates:
Received: 1 Sep. 2020 Revised: 27 May 2021 Accepted: 27 May 2021 Published: April 2021
Synopsis Tungsten carbide is a fine grey powder. It can be formed into shapes by compacting with the addition of a binder. Spherical particles are generally preferred in additive manufacturing as they pack together more efficiently than non-spherical particles, promoting a uniform powder bed density, better flowability, and elimination of internal cavities and fractures, resulting in a better quality of final product. The particle shape of powders can be transformed into spherical through the process of spheroidization. However, due to its high melting point, tungsten carbide could be difficult to spheroidize. Tungsten carbide was spheroidized using an inductively coupled radio frequency plasma at various plate powers between 9 and 15 kW. The influence of additional H2 in the sheath gas on the chemical composition of the tungsten carbide product was also investigated by means of XRD, which indicated that WC is converted to W2C at higher H2 concentrations. Optical analysis of SEM micrographs indicated that the spheroidization ratio increased with increased plasma energy. Keywords spheroidization, induction plasma melting, tungsten carbide.
How to cite:
Dire, R.M., Bissett, H., Delport, D., and Premlall, K. 2021 Evaluation of spheroidized tungsten carbide powder produced by induction plasma melting. Journal of the Southern African Institute of Mining and Metallurgy, vol. 121, no. 4, pp. 175–180. DOI ID: http://dx.doi.org/10.17159/24119717/1352/2021 ORCID R.M. Dire https://orchid.org/0000-00033107-3571 H. Bissett https://orchid.org/0000-00031034-7801
Introduction Tungsten carbide is a chemical compound used for various high-temperature applications. In its most basic form, tungsten carbide is a fine grey powder, which can be compacted with the addition of a binder (Co) and formed into shapes by a sintering process for use in industrial machinery, cutting tools, abrasives, armour-piercing shells, and jewellery (Pohanish, 2012). Carbides are one of the most interesting groups of materials in that they exhibit unique properties such as high chemical stability, low thermal expansion, high thermal conductivity, and high temperature resistance. Among the several carbide synthesis methods, carbothermic reduction is the most preferred technique in industry, as it allows the use of a wide range of materials as precursor (Devečerski, 2011). Carbothermic reaction involves the reduction of substances, often metal oxides (MOX), using carbon as the reducing agent (Earnshaw, and Greenwood, 1997) and reducing metal oxide nanoparticles to metallic elements in the presence of carbon by heat treatment in an inert atmosphere. The metallic element then reacts with the excess carbon to form a carbide (Barker, Swoyer, and Saidi, 2003). Tungsten carbide is extremely hard, and its surface resists impacts, scratches, and abrasion. It is almost as dense as gold, so tungsten carbide rings have a satisfying ‘heft’ when worn (Haynes, 2011). Some of the objects made from tungsten carbide can be manufactured through the additive manufacturing (AM) process, which is a fast-emerging technology in which a shape is fabricated using layer-by-layer deposition of a material in a bottom-up manufacturing method, allowing for the fabrication of three-dimensional (3D) parts with complex geometrical features that are difficult to manufacture using traditional machining techniques (Guo and Leu, 2013). The most important powder characteristics to consider for AM include: particle size, shape, density, porosity, surface area, and topography. Spherical powders are generally preferred as they pack together more efficiently than non-spherical particles, giving uniform powder bed density, better flowability (do not clog machinery), and result in higher quality of final products. The particle shape of powders can be transformed into spherical by the process of spheroidization. Spheroidization refers to heat treatment and/or a modification process that is used to convert granular shapes to spheroidal shapes (Hillert, 1962).
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Evaluation of spheroidized tungsten carbide powder produced by induction plasma melting There is a need in industry to convert angular shaped powders into spherical shaped powders. The following benefits would result from the spheroidization process: improved powder flowability, increased packing density, elimination of internal cavities and fractures, changing the surface morphology of the particles, and increasing the chemical purity (Hillert, 1962). Entities such as the National Laser Centre (NLC) at the CSIR uses tungsten carbide for thermal coating on various materials. Methods such as thermal spraying require spherical particles to ensure effective splat formation on the substrate or surface being coated. The South African Nuclear Energy Corporation SOC Ltd. (Necsa) has obtained a 15 kW inductively coupled radiofrequency (RF) thermal plasma system from Tekna Plasma Systems Inc., which is used for the spheroidization of various powders. The focus of this study is the spheroidization of irregularly-shaped pure tungsten carbide powder to investigate the capabilities of the plasma system. Spheroidization of powders should result in a spheroidization ratio as close as possible to 100%, while minimizing the fraction of material evaporated.
Radio frequency (RF) plasma spheroidisation
Plasma is classified as the fourth state of matter. Thermal plasmas can have very high temperatures based on the electron temperatures or densities. The temperature of the plasma is determined by the kinetic energies of the heavy particles such as the ions, atoms, or molecules (Haehn et al., 1986). Currently, methods such as inductively coupled RF plasma spheroidization and water atomization can be used to form spherical particles. However, most of the methods are not able to produce spherical tungsten carbide powders due to its high melting point (2870°C) (Ku et al., 2019). The RF plasma method is an exception. This process can be used to prepare powders
having high sphericity with minimum contamination, as the materials are not in contact with refractories and no impurities are added as a result of electrode evaporation (Kwak, 2010; Tong, 2015). Plasma spheroidization can manufacture spherical powder particles which have a uniform composition, high sphericity, and excellent flowability (Lu, Zhu, and Zhang, 2012). A schematic representation of the Tekna 15 kW induction plasma system is shown in Figure 1. This particular RF plasma system uses a PL-35M induction torch which can operate between 2 and 5 MHz (radio-frequency range). In this study the PL-35M torch was mounted on a reactor chamber equipped with a ’catch pot’ for the collection of the spheroidized tungsten carbide powder.The irregular-shaped feed powder was introduced into the plasma torch through a water-cooled probe with the assistance of an inert carrier gas (Ar). Inside the plasma ’flame’ the particles were rapidly melted, followed by rapid cooling to form spherical particles. Fine particles were formed due to evaporation, followed by condensation. These particles were collected in the cyclone and filter sections of the system, as shown in Figure 1. The plasma gases used were Ar or Ar/H2.
Experimental The tungsten carbide powder used in the experiment was a mixture of WC and W2C. The as-received powder was sieved into various fractions using a sieve shaker. Only the sieve fraction 45–75 μm was used in this study.
Spheroidization ratio, fraction of evaporation, particle density, and particle size distribution The powder was plasma treated at various conditions as indicated in Table I. For all experiments the central and sheath gas flow rates were kept constant. The powder feed rate was
Figure 1—Schematic representation of the 15 kW Tekna spheroidization system
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Evaluation of spheroidized tungsten carbide powder produced by induction plasma melting approximately 0.7 kg/h, reactor pressure was 85 kPa (abs.), and the carrier gas (Ar) flow rate was 2 standard litres per minute (slpm). After spheroidization, the collected powders were weighed in order to determine the evaporated fraction of the powder (condensed as very fine particles <<150 nm). SEM analyses of the densified powders were performed using a Quanta FEI 200 D SEM system. Image processing was applied on the backscattered electron (BSE) images. The Carl Zeiss Zen 2 Core software package was used to determine whether there was a quantifiable change in the particle morphology under different conditions. The average circularity was determined by evaluating a number of particles and making use of Equation [1]. The shape factor fcirc is the circularity, a function of the perimeter P and the area A: [1] In this instance the circularity varies between unity and nearzero, with the circularity of a circle equal to 1. Spheroidization should result in an increased circularity when compared to the irregularly shaped feed powder. The feed and the plasma-treated powders were also characterized according to density, using an AccuPyc II 1340 gas displacement helium pycnometer. The 10 cm3 sample cup was used for density analysis of three samples to determine the average density. Particle size distributions (PSDs) of all powders were determined using a Saturn DigiSizer II analyser.
Table I
hermal plasma conditions for the treatment of T 45–75 μm WC powder Plasma plate power (kW)
Hydrogen concentration in plasma (% v/v)
Energy consumption (kW.h/kg)
0.00 0.04 0.08
10.98 13.58 20.00
9 11 15
Phase composition Phase composition of the powders was determined by X-ray diffraction (XRD) using a Bruker D8 Advance diffractometer equipped with a Cu Kα (λ = 0.15418 nm) radiation source. Data was collected using a LynxEye position-sensitive detector. The phases were identified using ICDD's PDF4+ 2020 version database. The phase compositions were determined by Rietveld refinement.
Results and discussion Figure 2 shows the SEM BSE images of the feed and plasmatreated powders. The feed powder was irregularly-shaped and spheroidization of the powder had occurred at higher (11 and 15 kW) plasma plate powers and higher H2 concentrations (0.04 and 0.08% v/v).
Figure 2—SEM-BSE images at 155× magnification of the (a) tungsten carbide 45–75 µm feed powder and powders plasma-treated at (b) 9 kW, (c) 11 kW, and (d) 15 kW The Journal of the Southern African Institute of Mining and Metallurgy
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Evaluation of spheroidized tungsten carbide powder produced by induction plasma melting Table II
verage spheroidization ratios as determined by image processing software, with spherical particles having circularity A values of 0.85, and fraction of powder evaporation Plasma plate power (kW) As received 9 11 15
Spherical particles number
Irregular particles number
67 12 582 922
648 233 51 59
[2] . where P is the minimum energy per unit time required (W), m is the powder feed rate (g/s), cp is the specific heat capacity (J/K.g), Tm is the melting point, T0 is the room temperature (K), and Hm is the latent heat of fusion (J/g) (Jiang, and Boulos, 2006) of the powder. Therefore, for a given power input, a specific number of particles will be melted. When more power is applied, more particles can be spheroidized, as shown in Table II. This is due to the higher temperature of the plasma. The plasma, however, has a temperature profile and therefore some particles moving through the high-temperature regions will be spheroidized, while other particles might only be partially melted and do not completely spheroidize. APRIL 2021
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Fraction of evaporation (%)
9.4 4.9 91.9 94.0
The quantified changes in particle morphology as determined by image processing, as well as the fraction of evaporation and powder densities, are indicated in Table II. The fraction of evaporation for all experiments was low (< 4%), which can be attributed to the high boiling point (6 000°C) of tungsten carbide (Gignard, 1998). The densities of the plasma-treated powders were slightly higher than the feed powder due to densification of the particles through melting and void elimination, but also due to possible loss of carbon and conversion of WC to W2C. The spheroidization ratio of the feed powder was 9.4 %, indicating that some of the irregularly-shaped particles in the feed had satisfied the selected circularity value of 0.85 for spherical particles. The treatment at 9 kW resulted in a decreased spheroidization ratio. In this instance, melting of the particle surfaces, partial melting of the particles, and sintering of fine and larger particles occurred. Once these particles exit the plasma tail flame, rapid quenching of the oval and sintered particles occurs, resulting in a decreased spheroidization ratio compared to the feed powder. The sintering of fine and larger particles resulted in an overall increase in particle size, shown in Figure 3 and Table III. A bimodal PSD was observed after treatment at 9 kW, resulting in an increased mean and median compared to the feed and other plasma-treated powders. The feed powder had a unimodal distribution with a median of 58 µm, while the medians of the powder treated at 11 kW and 15 kW shifted slightly to a smaller particle size due to the spheroidization of the particles. The appearance of another, less significant, peak for both the 11 kW and 15 kW treatments might be due to the measurement of finer particles which affect the tail of the distribution. The spheroidization ratios obtained after treatment of the powder at 11 kW and 15 kW were 91.9% and 94.0% respectively. A minimum energy transfer per unit time was required, which results in the melting of powder at a given powder feed rate, given by Equation [2]:
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Spheroidization ratio (%)
– 2.67 3.19 0.90
Powder density (g/cm3) 16.414 16.756 16.750 16.743
Increasing the H2 concentration will also influence the spheroidization ratio. The thermal conductivity of Ar is much lower than that of H2 – 71 mW/m.K at 2 100 K (Chen and Saxena, 1975) compared with 1178 m.W/m.K for H2 (Blais and Mann, 1960). This results in more effective heat transfer from the plasma to the particles when using H2.
Phase composition The XRD spectra of the tungsten carbide powders before and after the RF plasma spheroidization are shown in Figure 4. There are eight major diffraction peaks with 2θ values of 39°, 47°, 51°, 61°, 69°, 74°, 106°, and 111°. These results conform to diffraction peaks of standard tungsten carbide powder with a body-centred cubic (bcc) lattice structure. There were two phases which were identified in all samples; WC and W2C. A minor peak at 80° indicated the presence of Co, which is added as a binder in the tungsten carbide powder. Table IV shows that the feed powder comprised both WC and W2C phases. The amount of W2C phase increased with plasma plate power and with increased H2 concentration. According to Gignard (1998), the calculated equilibrium composition of the C-H-W system shows that W2C is
Figure 3—Incremental volume percentage as a function of particle size for tungsten carbide feed powder treated at 9 kW, 11 kW, and 15 kW
Table III
SD of all powders as determined by light scattering P analysis
Sample Feed 9 kW 11 kW 15 kW
H2 (%v/v)
Mean (μm)
Median (d50) (μm)
Standard deviation (μm)
– 0.00 0.04 0.08
62.597 70.115 74.270 75.958
58.222 62.752 55.156 53.359
1.012 0.946 0.172 0.187
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Evaluation of spheroidized tungsten carbide powder produced by induction plasma melting
Figure 4—XRD spectra of the tungsten carbide feed and plasma-treated powders
References
Table IV
P roposed phases proportions of the feed and plasma treated powders as determined from XRD analysis
Jiang, X.L. and Boulos, M. 2006. Induction plasma spheroidization of tungsten and molybdenum powders. Transactions of Nonferrous Metals Society of China, vol.16. pp. 13–17.
Sample
Lu, X., Zhu, L., and Zhang, B. 2012. Simulation of flow field and particle trajectory of radio frequency inductively coupled plasma spheroidization. Computationalr Materials Science, vol. 65. pp. 13–18.
Feed 9 kW 11 kW 15 kW
H2
Phase composition (%)
(% v/v)
WC
W2C
– 0.00 0.04 0.08
35.6 21.4 14.3 14.4
64.4 78.6 85.7 85.6
more stable at higher temperatures than WC. This explains why a fraction of the WC converted to W2C during the spheroidization process and the W2C fraction increased with increasing plasma temperature.
Conclusion Tungsten carbide powder was successfully spheroidized utilizing an inductively coupled radio-frequency plasma process, using a reactor power greater than 9 kW with the addition of H2 to the sheath gas. Spheroidization increased the density of the powder. Spheroidization ratios greater than 90% were achieved at 11 kW and higher, with minimal powder evaporation occurring. The XRD results showed increased proportions of W2C when the power of the RF plasma was increased. Thus the power of the plasma had a direct influence on the formation of the W2C phase due to the reduction in WC content. The W2C proportion increased with plasma plate power and with increasing H2 concentration in the sheath gas. An increase in the W2C phase at higher plasma temperatures was due to the fact that W2C is more stable than WC at higher temperatures, as confirmed by the calculated equilibrium composition of the C-H-W system (Gignard, 1998).
Acknowledgements The authors would like to acknowledge the Nuclear Materials Development Network (NMDN) of the Advanced Materials Initiative (AMI), funded by the Department of Science and Innovation (DSI), for the financial support in conducting this study. The South African Nuclear Energy Corporation (Necsa) is acknowledged for their financial support. The following Necsa personnel are thanked for their contributions, S.J. Lotter and T.P. Ntsoane for the SEM and XRD analysis respectively, and M.M. Makhofane and P.C. Smith for the plasma experiments. M. Theron from the CSIR National Laser Centre is also thanked for supplying the tungsten carbide powder which was used for experiments. The Journal of the Southern African Institute of Mining and Metallurgy
Barker, J., Swoyer, M, and Saidi, J. 2003. A carbothermal reduction method for the preparation of electroactive materials for lithium ion applications. Journal of the Electrochemical Society, vol. 150, no. 6. doi: 10.1149/1.1568936 Blais, N.C. and Mann, J.B. 1960. Thermal conductivity of helium and hydrogen at high temperatures. Journal of Chemical Physics, vol. 32. pp. 1459–1465. Chen, S.H.P. and Saxena, S.C. 1975. Thermal conductivity of argon in the temperature range 350 to 2500 K. Molecular Physics, vol. 29. pp. 455–466. Devečerski, A. 2011. SiC synthesis using domestic resources. Process and Application of Ceramics, vol. 5. pp. 63–67. Earnshaw, N. and Greenwood, A. 1997. Chemistry of the Elements. ButterworthHeinemann, Oxford. p. 308. Gignard, N.M. 1998. Experimental optimization of the spheroidization of metallic and ceramic powders with induction plasma. Thesis. National Library of Canada, Sherbrooke, Quebec, Canada. Guo, M and Leu, N. 2013. Additive manufacturing: technology applications and research needs. Frontiers of Mechanical Engineering, vol. 8, no. 3. pp. 215–243. Haehn, R., Luederitz, E., Sattelberger, S., and Retelsdorf, H. 1986. New process for the production of cast tungsten carbide. Metal Powder Report, vol. 41, no.12. pp. 887–890. Haynes W. 2011. CRC Handbook of Chemistry and Physics. CRC Press, Boca Raton, FL. p. 496. Hillert, M. 1962. The formation of pearlite. Proceedings of Decomposition of Austenite by Diffusional Processes. Philadelphia, Pennsylvania, 19 October, 1960. Zackay, V.F. and Aaronson, H.I. (eds.). Interscience Publishers, New York, 1962. pp. 197–249. Ku, N., Pittari, J.J., Kilczewski, S., and Kudzal, A. 2019. Additive manufacturing of cemented tungsten carbide with a cobalt-free alloy binder by selective laser melting for high hardness applications. JOM, vol. 71, no. 4. pp. 1535–1542. Kwak, S. 2010. A review of switch mode sustain drivers with resonant networks for plasma display panels. IEEE Transactions on Industrial Electronics, vol. 57, no. 5. pp. 1624–1634. Pohanish, R. 2012. Sittig's Handbook of Toxic and Hazardous Chemicals and Carcinogens. (6th edn). Elsevier, Norwich. p. 2670. Tong, J. 2015. Fabrication of micro-fine spherical high Nb containing TiAl alloy powder based on reaction synthesis and RF plasma spheroidization. Powder Technology, vol. 28. pp. 9–15. u VOLUME 121
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1–2 July 2021 — The Mine Waste & Tailings Stewardship Conference 2021 Brisbane, Australia Website: https://www.ausimm.com/conferences-and-events/mine-waste-and-tailings/ 13–16 July 2021 — Copper Cobalt Africa Incorporating The 10th Southern African Base Metals Conference Avani Victoria Falls Resort, Livingstone, Zambia Contact: Camielah Jardine Tel: +27 11 834-1273/7 Fax: +27 11 838-5923/833-8156 E-mail: camielah@saimm.co.za Website: http://www.saimm.co.za 28 July – 22 September 2021 — 5th Mineral Project Valuation Hybrid Colloquium The Canvas Riversands, Fourways, Johannesburg Contact: Gugu Charlie Tel: +27 11 834-1273/7 Fax: +27 11 838-5923/833-8156 E-mail: gugu@saimm.co.za Website: http://www.saimm.co.za 3–4 August 2021 — DIMI (Diversity and Inclusion in the Minerals Industry) Hybrid Conference 2021 ‘Empowering the African minerals industry through diversity and inclusion’ The Canvas, Riversands, Fourways, South Africa Contact: Camielah Jardine Tel: +27 11 834-1273/7 Fax: +27 11 838-5923/833-8156 E-mail: camielah@saimm.co.za Website: http://www.saimm.co.za 16–17 August 2021 — Worldgold Hybrid Conference 2021 Misty Hills Conference Centre, Muldersdrift, Johannesburg, South Africa Contact: Camielah Jardine Tel: +27 11 834-1273/7 Fax: +27 11 838-5923/833-8156 E-mail: camielah@saimm.co.za Website: http://www.saimm.co.za 29 August–2 September 2021 — APCOM 2021 Minerals Industry Hybrid Conference ‘The next digital transformation in mining’ Misty Hills Conference Centre, Muldersdrift, Johannesburg, South Africa Contact: Camielah Jardine Tel: +27 11 834-1273/7 Fax: +27 11 838-5923/833-8156 E-mail: camielah@saimm.co.za Website: http://www.saimm.co.za
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NATIONAL & INTERNATIONAL ACTIVITIES 21–22 September 2021 — 5th Young Professionals Hybrid Conference 2021 ‘A Showcase of Emerging Research and Innovation in the Minerals Industry’ The Canvas, Riversands, Fourways, South Africa Contact: Gugu Charlie Tel: +27 11 834-1273/7 Fax: +27 11 838-5923/833-8156 E-mail: gugu@saimm.co.za Website: http://www.saimm.co.za 23–24 September 2021 — 3rd Global Engineering Online Symposium ‘Bridging the gap between research (academia) and industry while transitioning into 4IR’ Contact: Gugu Charlie Tel: +27 11 834-1273/7 Fax: +27 11 838-5923/833-8156 E-mail: gugu@saimm.co.za Website: http://www.saimm.co.za 26–29 September 2021 — The 16th International Ferroalloys Congress (INFACON XVI) Clarion Hotel & Congress, Trondheim, Norway infacon2021@videre.ntnu.no 18–19 October 2021 — Southern African Rare Earths International Conference 2021 Misty Hills Conference Venue, Muldersdrift, Johannesburg, South Africa Contact: Camielah Jardine Tel: +27 11 834-1273/7 Fax: +27 11 838-5923/833-8156 E-mail: camielah@saimm.co.za Website: http://www.saimm.co.za 26–27 October 2021 — SAMCODES Conference 2021 ‘Good Practice and Lessons Learnt’ The Canvas Riversands, Fourways, South Africa Contact: Gugu Charlie Tel: +27 11 834-1273/7 Fax: +27 11 838-5923/833-8156 E-mail: gugu@saimm.co.za Website: http://www.saimm.co.za 15–17 November 2021 — Global Tailings Standards and Opportuniries Hybrid Conference 2021 ‘For the Mine of the Future’ Misty Hills Conference Venue, Muldersdrift, Johannesburg, South Africa Contact: Gugu Charlie Tel: +27 11 834-1273/7 Fax: +27 11 838-5923/833-8156 E-mail: gugu@saimm.co.za Website: http://www.saimm.co.za 11 November 2021 — 17TH Annual Student Colloquium 2021 The Canvas Riversands, Fourways, South Africa Contact: Gugu Charlie Tel: +27 11 834-1273/7 Fax: +27 11 838-5923/833-8156
The Journal of the Southern African Institute of Mining and Metallurgy
E-mail: gugu@saimm.co.za Website: http://www.saimm.co.za
2022 16–19 May 2022 — 8th Sulphur and Sulphuric Acid Conference |2022 The Vineyard Hotel, Newlands, Cape Town, South Africa Contact: Camielah Jardine Tel: +27 11 834-1273/7 Fax: +27 11 838-5923/833-8156 E-mail: camielah@saimm.co.za Website: http://www.saimm.co.za 14–16 June 2022 — Base Metals 2022 Copper Cobalt Africa in association with the 10th Southern African Base Metals Conference Avani Victoria Falls Resort, Livingstone, Zambia Contact: Camielah Jardine Tel: +27 11 834-1273/7 Fax: +27 11 838-5923/833-8156 E-mail: camielah@saimm.co.za Website: http://www.saimm.co.za 4-8 July 2022 — SOMP Society of Mining Professors 32nd SOMP Annual Meeting and Conference Windhoek Country Club & Resort, Windhoek, Namibia Contact: Camielah Jardine Tel: +27 11 834-1273/7 Fax: +27 11 838-5923/833-8156 E-mail: camielah@saimm.co.za Website: http://www.saimm.co.za 9–13 July 2022 — SDIMI Sustainable Development in the Minerals Industry 10th International Conference Swakopmund Hotel and Entertainment Centre, Swakopmund, Namibia Contact: Gugu Charlie Tel: +27 11 834-1273/7 Fax: +27 11 838-5923/833-8156 E-mail: gugu@saimm.co.za Website: http://www.saimm.co.za 6–8 September 2022 — PGM The 8th International Conference 2021 Sun City, Rustenburg, South Africa Contact: Camielah Jardine Tel: +27 11 834-1273/7 Fax: +27 11 838-5923/833-8156 E-mail: camielah@saimm.co.za Website: http://www.saimm.co.za 7–8 September 2022 — THANOS International Conference on Enhanced Use of Thermodynamic Data in Pyrometallurgy Teaching and Research Mintek, Randburg Contact: Camielah Jardine Tel: +27 11 834-1273/7 Fax: +27 11 838-5923/833-8156 E-mail: camielah@saimm.co.za Website: http://www.saimm.co.za VOLUME 121
APRIL 2021
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Company affiliates The following organizations have been admitted to the Institute as Company Affiliates 3M South Africa (Pty) Limited
MSA Group (Pty) Ltd
AECOM SA (Pty) Ltd
Multotec (Pty) Ltd
Expectra 2004 (Pty) Ltd Exxaro Coal (Pty) Ltd Exxaro Resources Limited AEL Mining Services Limited Filtaquip (Pty) Ltd African Pegmatite (Pty) Ltd FLSmidth Minerals (Pty) Ltd Air Liquide (Pty) Ltd Fluor Daniel SA ( Pty) Ltd Alexander Proudfoot Africa (Pty) Ltd Franki Africa (Pty) Ltd-JHB AMEC Foster Wheeler Fraser Alexander (Pty) Ltd AMIRA International Africa (Pty) Ltd G H H Mining Machines (Pty) Ltd ANDRITZ Delkor(pty) Ltd Geobrugg Southern Africa (Pty) Ltd Anglo Operations Proprietary Limited Glencore Anglogold Ashanti Ltd Gravitas Minerals (Pty) Ltd Arcus Gibb (Pty) Ltd Hall Core Drilling (Pty) Ltd ASPASA Hatch (Pty) Ltd Aurecon South Africa (Pty) Ltd Herrenknecht AG Aveng Engineering HPE Hydro Power Equipment (Pty) Ltd Aveng Mining Shafts and Underground Immersive Technologies Axiom Chemlab Supplies (Pty) Ltd IMS Engineering (Pty) Ltd Axis House (Pty) Ltd Ingwenya Mineral Processing (Pty) Ltd Bafokeng Rasimone Platinum Mine Ivanhoe Mines SA Barloworld Equipment -Mining Joy Global Inc.(Africa) BASF Holdings SA (Pty) Ltd Kudumane Manganese Resources BCL Limited Leica Geosystems (Pty) Ltd Becker Mining (Pty) Ltd Longyear South Africa (Pty) Ltd BedRock Mining Support (Pty) Ltd Lull Storm Trading (Pty) Ltd BHP Billiton Energy Coal SA Ltd Maccaferri SA (Pty) Ltd Blue Cube Systems (Pty) Ltd Magnetech (Pty) Ltd Bluhm Burton Engineering (Pty) Ltd MAGOTTEAUX (PTY) LTD Bond Equipment (Pty) Ltd Malvern Panalytical (Pty) Ltd Bouygues Travaux Publics Maptek (Pty) Ltd Castle Lead Works Maxam Dantex (Pty) Ltd CDM Group MBE Minerals SA Pty Ltd CGG Services SA MCC Contracts (Pty) Ltd Coalmin Process Technologies CC MD Mineral Technologies SA (Pty) Ltd Concor Opencast Mining MDM Technical Africa (Pty) Ltd Concor Technicrete Metalock Engineering RSA (Pty) Ltd Council for Geoscience Library Metorex Limited CRONIMET Mining Processing SA (Pty) Metso Minerals (South Africa) Pty Ltd Ltd Micromine Africa (Pty) Ltd CSIR Natural Resources and the Environ- MineARC South Africa (Pty) Ltd ment (NRE) Minerals Council of South Africa Data Mine SA Minerals Operations Executive (Pty) Ltd Digby Wells and Associates MineRP Holding (Pty) Ltd DRA Mineral Projects (Pty) Ltd Mining Projections Concepts DTP Mining - Bouygues Construction Mintek Duraset MIP Process Technologies (Pty) Ltd Elbroc Mining Products (Pty) Ltd MLB Investment CC eThekwini Municipality Modular Mining Systems Africa (Pty) Ex Mente Technologies (Pty) Ltd Ltd
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APRIL 2021
VOLUME 121
Murray and Roberts Cementation Nalco Africa (Pty) Ltd Namakwa Sands (Pty) Ltd Ncamiso Trading (Pty) Ltd New Concept Mining (Pty) Ltd Northam Platinum Ltd - Zondereinde Opermin Operational Excellence OPTRON (Pty) Ltd Paterson & Cooke Consulting Engineers (Pty) Ltd Perkinelmer Polysius A Division Of Thyssenkrupp Industrial Sol Precious Metals Refiners Rams Mining Technologies Rand Refinery Limited Redpath Mining (South Africa) (Pty) Ltd Rocbolt Technologies Rosond (Pty) Ltd Royal Bafokeng Platinum Roytec Global (Pty) Ltd RungePincockMinarco Limited Rustenburg Platinum Mines Limited Salene Mining (Pty) Ltd Sandvik Mining and Construction Delmas (Pty) Ltd Sandvik Mining and Construction RSA (Pty) Ltd SANIRE Schauenburg (Pty) Ltd Sebilo Resources (Pty) Ltd SENET (Pty) Ltd Senmin International (Pty) Ltd SISA Inspection (Pty) Ltd Smec South Africa Sound Mining Solution (Pty) Ltd SRK Consulting SA (Pty) Ltd Time Mining and Processing (Pty) Ltd Timrite Pty Ltd Tomra (Pty) Ltd Ukwazi Mining Solutions (Pty) Ltd Umgeni Water Webber Wentzel Weir Minerals Africa Welding Alloys South Africa Worley
The Journal of the Southern African Institute of Mining and Metallurgy
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8th
SULPHUR AND SULPHURIC ACID CONFERENCE | 2022 16 MAY 2022 - WORKSHOP
Sulfuric Acid Catalysis - Key Parameters to Increase Efficiency and Lower Costs
17-18 MAY 2022 - CONFERENCE 19 MAY 2022 - TECHNICAL VISIT THE VINEYARD HOTEL, NEWLANDS, CAPE TOWN, SOUTH AFRICA
BACKGROUND The production of SO2 and sulphuric acid remains a pertinent topic in the Southern African mining and metallurgical industry, especially in view of the strong demand for, and increasing prices of, vital base metals such as cobalt and copper. The electric car revolution is well underway and demand for cobalt is rocketing. New sulphuric acid plants are being built, comprising both smelters and sulphur burners, as the demand for metals increases. However, these projects take time to plan and construct, and in the interim sulphuric acid is being sourced from far afield, sometimes more than 2000 km away from the place that it is required. The need for sulphuric acid ‘sinks’ such as phosphate fertilizer plants is also becoming apparent. All of the above factors create both opportunities and issues and supply chain challenges. To ensure that you stay abreast of developments in the industry, the Southern African Institute of Mining and Metallurgy invites you to participate in a conference on the production, utilization, safe transportation and conversion of sulphur, sulphuric acid, and SO2 abatement in metallurgical and other processes, to be held in 19 May 2022 in Cape Town.
FORMAT OF THE EVENT At this point in time, the event is planned as a full contact conference with international participation through web links. It is also planned to hold technical visits to nearby facilities. The situation will be constantly reviewed, and if it appears that the effects of the pandemic are still such as to pose a threat to the health and safety of delegates, this will be changed to a digital event.
OBJECTIVES •
•
•
•
To expose delegates to issues relating to the generation and handling of sulphur, sulphuric acid, and SO2 abatement in the metallurgical and other industries. Provide an opportunity to producers and consumers of sulphur and sulphuric acid and related products to be introduced to new technologies and equipment in the field. Enable participants to share information about and experience in the application of such technologies. Provide an opportunity for role players in the industry to discuss common problems and their solutions.
WHO SHOULD ATTEND The Conference will be of value to: Metallurgical and chemical engineers working in the minerals and metals processing and chemical industries Metallurgical/chemical/plant management Project managers Research and development personnel Academics and students Technology providers and engineering firms Equipment and system providers Relevant legislators
EXHIBITION AND SPONSORSHIP There are a number of sponsorship opportunities available. Companies wishing to sponsor or exhibit should contact the Conference Co-ordinator.
FOR FURTHER INFORMATION, CONTACT: Camielah Jardine, Head of Conferencing E-mail: camielah@saimm.co.za Tel: +27 11 834-1273/7 Web: www.saimm.co.za
Transportation
WORKSHOP SPONSOR