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Vol. 5. Num. 3.
Pages 199-292 (July - September 2016)
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Vol. 5. Num. 3.
Pages 199-292 (July - September 2016)
Original Article
DOI: 10.1016/j.jmrt.2015.12.001
Open Access
Effect of synthetic graphite and activated charcoal addition on the mechanical, microstructure and wear properties of AZ 81 Mg alloys
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Ram Prabhu T
CEMILAC, Defence R&D Organization, Bangalore, India
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Table 1. Chemical composition and designation of the AZ 81 Mg alloys.
Table 2. Mechanical properties and grain size of the AZ 81 Mg alloys.
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Abstract

In the present study, the grain refinement effects of carbon on AZ 81 (Mg–8%Al) alloys were investigated through the addition of synthetic graphite and activated charcoal inoculants. The mechanical and dry sliding wear properties of AZ 81 (Mg–8%Al) alloys with and without grain refiners were studied. The composition, microstructure and wear surface of the cast alloys were analyzed with the aid of atomic absorption spectroscopy, optical and scanning electron microscopes. The wear test was conducted at a speed of 5.37m/s and two loading (20N and 30N) conditions in a pin on disc wear apparatus. The mean grain size of the alloy is significantly reduced from about 185 to 32μm when the carbon content is increased to 0.98% C. Both tensile strength and ductility increase with increasing carbon content. Also, the wear rate decreases with the increase of carbon content in the alloy. The activated charcoal with 0.98wt% C is found to be a better grain refiner in improving the properties of tensile behaviour and wear resistance. The fracture surface morphology shows the brittle intergranular fracture. Analyses of wear surface morphology show that abrasive and delamination wear mechanisms are responsible for the wear loss in the AZ 81 alloy.

Keywords:
Mg alloys
Wear
Microstructure
Mechanical properties
Scanning electron microscopy
Full Text
1Introduction

Mg alloys are widely used in the lightweight applications because of their several attractive properties such as very high strength to weight ratio, excellent damping capacity, good electro-magnetic shielding, high recycle ability and machinability [1,2]. Recently, their importance in aerospace and automobile industries is regaining significantly because of the implication of the huge cost saving obtained through the weight reduction and fuel economy, and the availability of more environment-friendly processing technology. Mg–Al based alloys are one of the primary choice for manufacturing gear box housing, accessory drives and other non critical stress parts in aircrafts, and the wheel rims, chassis and driving gears and so on in automobiles [3,4]. An excellent castability characteristic of Mg–Al alloys allows the fabrication of complex and thin section parts through conventional sand casting process [5]. In addition, low cost, and excellent damping capacity make this group of alloys as an ideal choice for the gear box housing of the aircraft and automotive applications. In spite of these benefits, the strength and ductility of this alloy is still a major concern due to the coarser grain structure and the limited slip systems available in a hcp crystal structure. Particularly, the grain structure is much coarser in the sand casting processed parts.

The grain refinement is the only way that improves the mutually exclusive strength and ductility properties in the metal alloys. Melt superheating, native grain refinement through the control of impurities level, Elfinal process (FeCl3), particle inoculations (carbonaceous gases bubbling, solid carbon additives, SiC, TiC, AlN, TiB2), master alloy additions (Al–Ti–B, Al–60% Mn), Ca or Sb or Sr additions, rapid solidification, severe plastic deformation (ECAP, ARB), melt conditioning ultrasonic treatment, electromagnetic stirring are the major grain refinement technologies in Mg alloys [3,5]. Among these technologies, particle inoculations are found to be very effective in Mg–Al alloys. Particularly, carbon based inoculants seem to show remarkably greater grain refinement due to the better nucleation potency. Also, it does not require higher super heat temperature, large melt volume, and short melt holding time. Thus, the processing time and the crucible wear are reduced by the carbon inoculation method [6]. However, it should be noted that some of the carbon inoculants (Nucleant 5000, Wax fluorspar-carbon) are currently proprietary [7,8]. Information about the carbon based inoculation is scarce in the open literature. Available reports are mostly based on the C2Cl6, C6Cl16, MgCO3, graphite powder, paraffin, carbonaceous gases (e.g. CO, CO2, CH4), MnCO3, CaC2, SiC additions in the grain refinement of AZ 91 alloys [2,6,9,10]. The main limitation of these reports is that they lack in providing clear information about the compositional limit of carbon. Particularly, the use of carbonaceous gases, especially chloride based, is prohibited because of the toxicity (dioxins formation) and green house gas emissions [11]. Studies on the grain refinement effects of solid carbon inoculation are also limited. Particularly, the effects of activated charcoal and synthetic graphite on grain refinement and mechanical properties of AZ 81 alloys have not been investigated so far.

The growing demand for lightweight brakes and engine sliding components (piston, cylinder bores) makes the researchers to focus on the sliding behaviour of Mg alloys [12,13]. In general, studies on wear behaviour of Mg alloys are limited. There have been some reports exploring the effects of load, sliding speed and temperature on the wear behaviour of AZ 91 alloys [12–17]. These reports have focused on understanding the wear mechanisms prevailed for various sliding speed/load/sliding distance combinations. Further, consolidation of these reports gives a wear map delineating the mild and severe wear regimes in the AZ 91 alloy. However, the wear testing condition in the above reports is limited to low speed (<3m/s) ranges. The wear reports on higher speed conditions are not available. In addition to that, the wear behaviour of Mg–8% Al has not been investigated so far. The grain size effects on the wear resistance of the Mg–Al alloys have been poorly understood.

In light of the above facts, an experimental study is designed with two objectives: (1) the effect of solid carbon addition (synthetic graphite and activated charcoal) on the grain refinement and mechanical behaviour of AZ 81 alloys and (2) the effect of grain refinement on the sliding wear behaviour of AZ 81 alloys at high speed (5.37m/s) and two load (20 and 30N) conditions.

2Materials and methods2.1Materials processing and testing

The AZ 81 Mg alloy (Mg/Al-8%/Zn-0.5–0.8%/Mn-0.2–0.4%) (wt%) was prepared by the proper melt assessment in a mild steel crucible installed in the electrical resistance furnace. The melting was conducted in air under the protective cover of proprietary fluxes supplied by Magnesium Electron, UK. Two types of grain refiners were used in the study: (1) synthetic graphite (size: 21–30μm) and (2) activated charcoal (0.04wt% and 0.98wt%) (size: 25–30μm). The grain refiner was packed up in an Al foil and plunged in the liquid metal. Four types of alloy melts, (1) alloy with synthetic graphite grain refiner, (2) alloy with activated charcoal of 0.04wt% grain refiner, (3) alloy with activated charcoal of 0.98wt%, and (4) alloy without grain refiner were prepared. The super heat and pouring temperature of the alloy were 730°C and 690°C, respectively. The liquid metal was poured in the sand casting rectangular mould. The sand mould was prepared using the air settling silica sand. The gating design used for the present study is a bottom pouring design with the unpressurized gating ratio of 1:4:4. The melt pouring was conducted under the sulphur protective atmosphere to avoid any possible oxidation or burning. The proprietary ceramic filters were placed in the runner and the ingates to filter inclusions (sand, oxides, and flux) carried by the molten metal.

After the casting, the alloys were solutionized in three steps to refine the β-Mg17Al12 precipitates:

(1) Solutionizing at 410±5°C for 6h, (2) solutionizing at 350±5°C for 2h, ±5min, (3) solutionizing at 410±5°C for 10h followed by normal air cooling. The heat treatment was carried out under the SO2 protective atmosphere created from the FeS and S mixture placed inside the furnace. As this alloy was not amenable to the precipitation hardening, the quenching operation and the subsequent ageing treatment were not given to the alloy.

The composition of the cast melts was analyzed using the atomic emission spectrometer (SPECTROVAC). The results of composition analysis were reported in Table 1. The microstructure samples were prepared using standard metallographic techniques. The etchant used to reveal the microstructure was 5% Nital. The microstructure of the alloys was characterized using optical (Nikon Epiphot) and scanning electron microscopes. The grain size was measured from the micrographs using the Image J freeware. The hardness was measured in a Brinell scale (WOLPERT hardness tester). The indentation load and the size of the indenter used were 100N and 1mm steel ball, respectively. An average of five readings was taken as a hardness of the alloy. The tensile strength of the alloys was measured at a cross head speed of 1mm/min in a 25kN Instron universal testing machine. Tensile tests were carried out according to the ASTM E8/E8M standard at 28°C. The schematic of a tensile test sample is shown in Fig. 1.

Table 1.

Chemical composition and designation of the AZ 81 Mg alloys.

Composition (wt%)  Code  Al  Zn  Mn  Be  Mg 
Alloy without grain refiner  7.95  0.65  0.28  –  0.0014  91.12 
Alloy with synthetic graphite  8.06  0.6  0.23  0.03  0.0013  91.08 
Alloy with activated charcoal  8.1  0.64  0.25  0.04  0.0015  90.97 
Alloy with activated charcoal  8.15  0.6  0.27  0.98  0.015  90 
Fig. 1.

Schematic of tensile test specimen.

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The alloys were tested in a pin on disc tribometer apparatus to evaluate their wear resistance. Tests were performed at ambient temperature (25°C) and 60% humidity under the dry condition. The wear test parameters are: sliding speed 5.37m/s (700rpm), load 20 and 30N. The ASTM G99 standard was used to conduct the wear test. A schematic illustration of a pin on disc tribometer apparatus is shown in Fig. 2. The cylindrical shape of AZ 81 alloys with the diameter and length of 6mm and 30mm, respectively, was used as a pin. The contact surface of the pins was polished to 1μm (mirror polish) and cleaned in an ultrasonic bath containing acetone solution before the test. The disc used was a 160mm diameter EN 32 hardened steel. The hardness and an average surface roughness (Ra) of the disc were 62 HRC and 1.6μm respectively. The wear track perimeter of the pin was 460.3mm. The wear test was conducted for a 10km sliding distance. After every 2.5km, the pin and the disc were thoroughly cleaned with acetone solution to remove the loosely attached wear debris deposits. The wear loss of the pin was measured in terms of mass (mg) in a digital balance (0.1mg precision) for every 2.5km. The volume loss (ΔV) was estimated by dividing the mass loss with the density. The volumetric wear rate was computed using the following relation:

where WR is the volumetric wear rate (mm3/m), ΔV is the volume of material worn out (mm3), and D is the sliding distance (m).

Fig. 2.

Schematic of the Pin-on-Disc apparatus illustrating the pin position, wear track, and disc rotation direction.

(0.11MB).

The surface and cross section of the wear tested samples were examined under the optical microscope to identify the prevailing wear mechanisms and to understand the severity of the wear condition.

3Results and discussion3.1Microstructure and mechanical properties

The micrographs of the AZ 81 alloys with and without grain refiners are shown in Figs. 3 and 4. The microstructure of the alloys shows distinct equiaxial α-Mg grains with preferentially segregated β-Mg17Al12 discontinuous intermetallic precipitates along the grain boundaries. Particularly, the alloy without grain refiner shows the large amount of cellular eutectic structure inside the grains near the segregated β-Mg17Al12 precipitates, as seen in Fig. 3(a). The observed cellular structure in the alloy is the result of slow cooling from the sand mould. The interface of the eutectic cellular structure is the preferable location for the impurities segregation. It is also prone to casting defects formation. These defects weaken the alloy. Thus, the cellular structure is undesirable in AZ 81 alloys because it reduces the strength and ductility of the alloy. In contrast, alloys with grain refiner do not show the extensive cellular structure. The carbon grain refiners change the solidification pattern and restrict the formation of cellular structure. The carbon particles are found to be very tiny and distributed uniformly in the microstructure. They are found mostly close to the grain boundaries. It is noted that the grain refinement by carbon inoculations is observed to be uniform in the microstructure. Another important observation in the microstructure is the presence of micro porosities in the β-Mg17Al12 precipitates. These pores are potential sites for crack nucleation during tensile loading. The addition of carbon does not seem to influence the thickness, amount and distribution pattern of β phases. There is no observation of dendritic structure in the cast alloy.

Fig. 3.

SEM images of microstructure of the AZ 81 alloys: (a) alloy without grain refiner, (b) alloy refined with synthetic graphite, (c) alloy refined with activated charcoal of 0.04%, and (d) alloy refined with activated charcoal of 0.98%. (1) Cellular structure, (2) discontinuous Mg17Al12 precipitates, (3) micro porosities and (4) carbon particles.

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Fig. 4.

Optical microstructure of the AZ 81 alloys showing the grain refinement with increasing carbon addition, (a) alloy without grain refiner, (b) alloy refined with synthetic graphite (C: 0.03%), (c) alloy refined with activated charcoal (C: 0.04%), and (d) alloy refined with activated charcoal (C: 0.98%).

(1.39MB).

The grain size measurement results show that the alloy without grain refiner has the mean grain size of 185μm. The addition of grain refiners shows a significant grain size reduction in AZ 81 alloys, as seen from Table 2. These results confirm that the carbon has a definite effect in grain refinement of AZ 81 alloys. Various studies on AZ 91 alloy report that the formations of binary and ternary Al4C3 and Al2MgC2, Al2CO carbide nucleant in the melt provide the large number of nucleation sites for the grain formation [2,5,18–20]. To understand the nucleants responsible for the grain refinement in the present study, it is important to note the super heat temperature. The super heat temperature for the present case is 730°C. A thermodynamic calculation by Wang et al. [21] showed that Al4C3 is the most potential nucleant below 740°C in Mg–Al alloys because of the strong negative Gibbs free energy value (−47.8kJ/mol) for the Al4C3 formation. Also, the calculation shows that the partial pressure of oxygen in the molten alloy is about 6×10−51atm at 740°C, which makes the possibility of Al2CO formation highly unlikely. Further, the crystal structure of α-Mg and Al4C3 is same and the lattice misfit (Δa=3.8% and Δc=4.2%) between them is very low. These crystallographic facts strongly suggest that the Al4C3 nucleants are responsible for the grain refinement of AZ 81 alloys. The carbon inoculants provide grain refinement of AZ 81 alloys by two ways: (1) the particle pins the grains and restricts their growth and (2) the carbon segregation at the advancing solid–liquid interface during solidification enhances the constitutional undercooling. The undercooling increases the nucleation ratio and also, hinders the growing grains [6,22].

Table 2.

Mechanical properties and grain size of the AZ 81 Mg alloys.

Code  Mean grain size (μm)  Hardness (BHN)  0.2% proof strength (MPa)  Tensile strength (MPa)  % Elongation  Density (×10−3g/mm3
185±24  74.7±4.5  79  80  2.5  1.77 
114±21  83.6±2.7  97  117  2.9  1.76 
106±29  84.8±1.6  99  131  4.5  1.76 
37±15  96.7±1.6  193  219  6.1  1.76 

It is important to note that the alloy refined with the activated charcoal of 0.98wt% has the finest mean grain size of 47μm. This size is 74% lesser than the unrefined alloy grain size. The present result of grain refinement by 0.98% activated charcoal is relatively better than the other reported results [22]. It is because the higher amount of carbon in the melt maximizes the available nucleants. These nucleants restrict the growth of grains by grain pinning effect that results in grain refinement. An Interdependence model by Greer et al. [23] states that the grain size of the as-cast alloys is inversely related to the availability of the nucleants. Higher amount of nucleants shrinks the nucleation free zone in the melt and enhances the grain refinement. Thus, the increase of carbon content promotes the grain refinement significantly by creating more nucleant sites.

Tensile properties and hardness of the Mg alloys with and without grain refiners are given in Table 2. From the results, it is understood that the grain refiners significantly improve the hardness, strength and ductility properties of the AZ 81 alloy. Between the synthetic graphite and the activated charcoal, the charcoal provides slightly better grain refining effects, and thus better tensile properties and hardness for the equal weight % of carbon. As the carbon content increases, the size of the grains is significantly reduced. The tensile properties and hardness of the alloy are improved substantially due to the better grain refinement. The improved ductility with an increase of the carbon content is partly attributed to the reduced cellular structure. It is also reported that carbon addition promotes the formation of Al5Mn5, Al8(Mn,Fe)5 particles that assist in grain refinement of Mg–Al alloys [24,25]. Besides restricting the grain growth, the particles by itself act as dispersoids. They impede the dislocations very effectively resulting in significant improvement of strength through the Orowan (dislocation–particle interaction) strengthening mechanism. The particle removed sites in the fractographs, as seen in Fig. 5, confirm the contribution of strengthening by particles. The absence of cracked or broken particles in the fractographs indicates that the crack would have propagated along the particles in an intergranular manner and removed the entire particles rather than the particle fracture.

Fig. 5.

Fractographs of the AZ 81 alloys: (a) alloy without grain refiner, (b) alloy refined with synthetic graphite (C: 0.03%), and (c and d) alloy refined with activated charcoal (C: 0.98%). (1) Particles removed sites, (2) micro porosities, and (3) crack propagation along the grains.

(1.56MB).

Fractographs of the AZ 81 alloys with and without grain refiners are shown in Fig. 5. Fracture features are similar for all the alloys. Cracks are found to initiate at two locations: (1) pores located in the β-Mg17Al12 phase and (2) α-Mg/β-Mg17Al12 interfaces. As the crystal structure of β phase (BCC) is incompatible with the α-Mg matrix (HCP) and also, the β phase is soft and has lower strength than the α-Mg grains, the interfaces are more vulnerable to crack initiation. These cracks propagate along the grain boundaries, as seen in Fig. 5, leading to the brittle intergranular fracture of the alloy. Tiny non-interconnecting voids and secondary microcracks are also observed in the fractographs. There is no evidence of plastic flow in the alloy.

3.2Wear behaviour

The wear rate of AZ 81 alloys as a function of sliding distance for two load conditions (20N and 30N) is shown in Figs. 6 and 7. In general, the wear rate curves show two distinct states: (1) transition state and (2) steady state. In 20N load condition, the wear rate rises sharply in the transition period and thereafter, it stabilizes in the steady state for all the cases. The stabilization of the wear rate is attributed to the formation of a mechanical mixed tribolayer. In the transition state, the slope of the wear rate curve is much steeper for the AZ 81 alloy without grain refiner (A) compared to the AZ 81 alloy with 0.98% C (D). In the steady state, the AZ 81 alloy with 0.98% C (D) shows the lowest and constant wear rate. In contrast, the wear rate of the AZ 81 alloy without grain refiner (A) increases with the sliding distance in both the transition and steady state and it is found to be the highest among all the cases. Particularly, the increase in the wear rate with the sliding distance in the steady state for the A sample is attributed to the third body abrasion. The third body abrasion wear is caused by the excessive amount of loose wear debris trapped at the contact interface. The wear rate of the alloys refined with synthetic graphite (B) and activated charcoal of 0.04wt% (C) is in between A and D types. Similar trends are observed for the 30N load condition except that the slope of the wear rate with the sliding distance in the steady state is relatively lesser for all the cases. The above results imply that the wear resistance of the alloy is directly proportional to the hardness and strength properties and inversely related to the grain size. Thus, the grain refinement plays a vital role in improving the wear resistance of the AZ 81 alloys. Constant wear rate after the transition period gives the cue that the friction temperature rise during the sliding process may not be sufficient enough to soften and/or coarsen β-Mg17Al12 precipitates or grains in the AZ 81 alloy.

Fig. 6.

Wear rate as a function of sliding distance for the load and speed condition of 20N and 5.37m/s respectively.

(0.16MB).
Fig. 7.

Wear rate as a function of sliding distance for the load and speed condition of 30N and 5.37m/s respectively.

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Comparing the results of Figs. 6 and 7, it is clearly seen that the wear rate of the AZ 81 alloys increases with an increase of the applied load. The slope of the curve is similar in the transition period for both load conditions. The magnitude of the wear rate is almost twice for any particular sliding distance for all the cases when the load increases from 20N to 30N. This hints that the wear rate may have a linear relation with the load. However, more experiments at higher loads are required to prove this argument. Similar linear trend between the volume loss and the sliding distance was reported for AZ 91 alloys [13]. However, the testing condition (5–222N and 0.1m/s) in their study is completely different from the present study. No abrupt change in the wear rate with an increase of the applied load indicates that the alloy has not yet reached the severe wear regime, and a thermal equilibrium is reached between the contact couples above the sliding distance of 2500m due to the steady state mild wear condition in the AZ 81 alloy [13].

The wear surface images of the AZ 81 alloy refined with 0.98% C are shown in Fig. 8. The counter surface asperities and the strain hardened loose metal debris at the contact interface plow the Mg alloy continuously. This forms the abrasive groove lines and ridges parallel to the sliding direction. These features are of typical abrasive wear. The plowing process dislodges the material to either side of the groove line and to the edge of the samples, as seen as the burrs of a large lump at the edge of the sample in Fig. 8. Thereafter, the material is removed by the micro ploughing and cutting processes. Surface cracks are also visible in the surface nearly perpendicular to the sliding direction. The abrasive groove lines are deeper with an increase of the applied load, as seen in Fig. 8(a). The cross section of the wear tested sample in the unetched and etched condition, and the wear debris morphology of the AZ 81 alloy refined with 0.98% C are shown in Fig. 9. Fig. 9(a) and (b) clearly shows many delamination cracks in the subsurface running towards the surface. This confirms the operation of a delamination wear in addition to the abrasive wear. The wear debris has a large, irregular flaky shape and looks similar to a lustrous metal in the naked eye, as seen in Fig. 9(c). This debris may have formed through the subsurface delamination wear mechanism. The etched subsurface shows no major change in the grain size at the subsurface. Also, there are no changes in the phase distribution or the microstructure. It implies that the friction temperature rise is not sufficient enough to cause the coarsening of the grain or β-Mg17Al12 precipitates at the sub surface. The observed elongated grains close to the surface indicate that the subsurface has undergone significant strain hardening due to the plastic deformation.

Fig. 8.

Worn top surface view of the Mg alloy refined with activated charcoal (C: 0.98%) (a) 20N tested, (b) 30N tested showing (1) deep scratches and abrasive grooves created by counter surface asperities, (2) burr formed by heavy plowing, and (3) surface cracks, Arrow indicates the sliding direction.

(0.6MB).
Fig. 9.

Sectional view of the 30N load wear tested Mg alloy refined with activated charcoal (C: 0.98%) (a and b) sub surface in unetched condition, (c) wear debris morphology, and (d) sub surface in etched condition. (1) Delamination cracks running to surface and (2) strain hardened region showing elongated grains.

(1.16MB).
4Conclusions

The grain refinement effects of solid carbon inoculation in AZ 81 alloys are investigated for various carbon contents. The tensile and wear properties are studied in detail and compared to identify the optimum amount of carbon. The main results of the present investigation are as follows:

  • The grain refinement, tensile and wear resistance properties of the AZ 81 alloys increase with an increase of the carbon content.

  • The mean grain size is reduced from about 185 to 32μm when the carbon content is increased to 0.98% C. The extent of cellular structure significantly reduces with an increase of the carbon addition. The alloy refined by activated charcoal of 0.98% C shows the highest strength, ductility and wear resistance properties.

  • Fractography studies show that the alloy with 0.98% C fails by the intergranular brittle fracture in tension.

  • The wear rate increases linearly with an increase of the applied load. The presence of groove lines in the wear surface and the sub surface cracks confirm that both the abrasive and the delamination wear mechanisms are responsible for the wear loss in the AZ 81 Mg alloy.

Conflicts of interest

The author declares no conflicts of interest.

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Copyright © 2015. Brazilian Metallurgical, Materials and Mining Association
Journal of Materials Research and Technology

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