Development of Si3N4-AlN-Y2O3composite cutting tool

J.V.C. Souza1, C. Santos,2, C. A. Kelly2, M. V.Ribeiro3, K.Strecker4 , O.M.M. Silva5

1INPE - Av. dos Astronautas,1.758, S. J. Campo s - SP, CEP. 12245-970, Brazil

2FAENQUIL-DEMAR – Polo Urbo Industrial, Gleba AI-6, s/n, Lorena - SP, CEP. 12600-000, Brazil

3FEG-UNESP – Av.Ariberto Ferreira da Cunha, 333, Guaratinguetá - SP, CEP. 12516-410, Brazil

4UFSJ - Praça Frei Orlando 170, São João del Rei - MG, CEP. 36307-352, Brazil

5CTA-IAE/AMR - Pça. Marechal do Ar Eduardo Gomes, 50, S. J. Campo s - SP, CEP. 12228-904, Brazil

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Abstract

In this work, a new composition of the Si3N4 ceramic cutting tools was developed and characterized. Powder mixture based on Si3N4-AlN-Y2O3 ternary system was developed aiming the obtaining of SiAlON ceramic, a solid solution of AlN and Y2O3 in the Si3N4 structure. Compacts were sintered at 19000C, and characterized by X-Ray Diffraction and Scanning Electron Microscopy. The Si3N4-AlN-Y2O3 ceramiccomposite cutting tool insert presented hardness of 21GPa and fracture toughness of 5.6MPam1/2. Furthermore, XRD patterns indicate the presence of α-SiAlON as predominant crystalline phase and elongated grains were observed in the SEM analysis. Machining studies have been conducted using the newly developed cutting tool insert on compacted graphite iron(CGI) workpiece.

Keywords:Ceramiccutting tools; Flank wear; Machining, Finish surface, Compacted graphite iron.

* Corresponding author. Tel.: +55 (12) 3945-6679, 39456989, FAX: +55 (12) 3945-6717

E-mail address: / (J.V.C. Souza).

1. Introduction

The machining is a fabrication technique where different materials are removed from a part using a tool with a small hard tip. In order to quickly fabricate parts, high-speed cutting is desired. These higher speeds, however, lead to a faster degradation of the tool tip, which requires more frequent replacements of the cutting edge. Ceramic cutting tools [1-3]play important roles in the development of advanced manufacture technology, and the increased application of ceramic cutting tools in proportion is inevitable. Types and brands of commonly used ceramic tools were summarized. Compared with traditional high speed steel and hard alloys, advanced ceramic cutting tool has the advantage of high hardness, chemical stability etc.

Silicon nitride (Si3N4) based ceramic tools are widely used as cutting tools because of their high hardness, high thermal conductivity and low thermal expansion coefficient that give better thermal shock resistance than other ceramic materials. Under high-speed machining conditions, the elevated temperatures developed in the cutting zone (tool-chip and tool-workpiece contact), enhance chemical and mechanical wear mechanisms[2-4].

SiAlON (or ’) is a solid solution of Si3N4 of composition MxSi12-(m+n)Al(m+n)OnN16-n, where M represents the metallic cation used for the stabilization of this phase and Si and N are partially substituted by Al and O, respectively, in the crystal structure. This ceramic material exhibits higher creep resistance, hardness and oxidation resistance than Si3N4 [5-7]. Its improved high-temperature properties are due mainly to the absence or reduced amount of intergranular secondary phases, because of the incorporation of the metallic cation of the liquid phase into the Si3N4 structure forming SiAlON. The metal ions used for the stabilization of the -Si3N4 structure are also responsible for the densification during the sintering, the improvement of the microstructural characteristics and the mechanical properties of these ceramic materials. Yttrium Y+3is one of the cations used to stabilize the ’-phase at higher temperatures, generating elongated grains and therefore increased fracture toughness.

The overall goal of the present study is to explore silicon nitride capabilities for machining of compacted graphite iron (CGI). To achieve this goal, the previously validated of development new cutting tools ceramic to apply on turning compacted graphite iron (CGI) and select better machining conditions. Machinability tests were conducted and the resulting data were analyzed, including tool wear rates, surface finish and temperature so that machining parameters which can extend tool life could be identified.

1.2. Properties of Compacted graphite iron

The material properties of compacted graphite iron (CGI) offer better strength and stiffness than gray iron, and better castability, machinability, and thermal conductivity than ductile iron, making it ideal for components that undergo both mechanical and thermal loading [8]. There is almost always some spheroidal graphite present in CGI, which is referred to as percent nodularity, which is ideally between 0 and 10 %, although 20 % is allowable depending on the application. Compacted graphite iron (CGI) tends to hold its strength and elongation properties well until approximately 400°C[9].

The CGI has a higher tensile strength and higher stiffness, it is difficult to machine. Additionally, the pearlitic CGI structure is more difficult to machine than ferritic CGI [9]. The properties of CGI necessitate that high torque and stiffness machine tools be used and Dawson[10] notes that 20–30 % higher spindle power is required. The greatest losses in tool life occur during turning and cylinder boring[11]. Dawson et al.[12] state that during low speed cutting using carbide tooling, a 50 % reduction in tool life is seen as compared to gray iron.

2. Experimental

Commercially available, high-purity α-Si3N4 powder, Y2O3 powder and AlN Grade-A Powder (H.C. Starck-Germany) were used as starting powders. Powder batch containing 82.86 wt.% ofα-Si3N4, 6.51%Y2O3and 10.63%AlN was obtained by ball-milling for 4h, in a plastic jar using Si3N4 balls and ethanol as the milling-media. The ball/powder weight ratio used was 5:1. Powder mixture was dried at 100°C for 24h and sieved. Green compacts were cold uniaxially pressed under a pressure of 50MPa. Prismatic green samples (16.4×16.4 × 8mm3) were encapsulated in plastic bags and isostatically pressed under pressure of 300MPa for 2min. The green relative density of the samples was obtained using the relation between the measured density (weight/volume) and theoretical density of the mixture.

Sintering was carried out in a graphite-resistance furnace (Thermal Technology 1000-4560-FP-20). The compacts were heated under vacuum, at 600°C-10 min. using a heating rate of 5°C/min. The low heating rate was employed in order to allow the organic binder to burn slowly and prevent the damage of the samples. From 600°C to a final sintering temperature (19000C-120min), the samples were heated using 15°C/min as rate. At 12000C, nitrogen atmosphere (0.1MPa) was used in the sintering and the cooling rate was 150C/min.

The relative density of the sintered samples was determined by the Archimedes’ method in distilled water. Phase analysis was done by X-ray diffraction (XRD) using a Phillips PW-1380/80 X-ray Diffractometter, Cu-Kα radiation (λ=1.5418Å), comparing the diffraction patterns with the JCPDS files. X-ray diffraction analysis was conducted on a cross section of the samples.

For microstructural analysis, the samples were ground, polished, and chemically etched by a 1:1 mixture of NaOH and KOH at 5000C for 10 min. Crack propagation and fracture surfaces of the specimens were observed using scanning electron microscope (LEO 1450VP) at an accelerating voltage of 20kV. A gold coating was applied on the surface of the etched samples in order to minimize the surface charging under electron beam.

The hardness was determined by Vicker’s indentations for 30 s under an applied load of 10 kgf. For statistical reasons, 21 indentations have been made per sample. The fracture toughness has been determined by the measurement of the crack length created by the Vicker’s indentations. The calculation of the fracture toughness, was done by the relation proposed in Equation (1) [13],valid for Palmqvist type cracks:

KIC = 0.16(E/H)1/2.F.b-3/2 (1)

Where: KIC = fracture toughness [MPa.m1/2]; E = Young modulus of material [GPa]; HV = Vickers hardness [GPa]; b = crack size [m] and F = applied load [N].

2.1. Machining test

All experiments were carried out on a computer numerical control (CNC) lathe (Romi, Mod. Centur 30D) under dry cutting condition. The sintered ceramic cutting tool were cut and ground to make SNGN120408 (12.7x12.7x4.76mm, 0.08mm nose radius and 0.2mm x 20° chamfer). A tool holder of CSRNR 2525M 12CEA type (offset shank with 15° [75°] side cutting edge angle, 0° insert normal clearance and 25 mm x 25 mm x 150mm) was used for the cutting experiments. The cutting performance of the ceramic cutting tool was tested by machining compacted graphite iron.

The cutting tests for machining of compacted graphite iron were performed at cutting speed of Vc1=200, Vc2=350 andVc3=500m/min with a feed rate of 0.20mm/rev and a depth of cut of 0.5mm. The compacted graphite iron presented tensile strength of 500MPa, elastic modulus of 140GPa, thermal conductivity of 35 W/m-K and 225BHN. Since the workpieces were manufactured by the rolling process in order to eliminate the problems encountered in the first pass, 1.5–2mm of skin was removed with a separate insert before the experiments. This application reduces the vibrations since it improves the roundness accuracy of the workpieces. The dimension of work material was about 60mm in intern diameter, external diameter 120 and 400mm in length. The wear of the tools was determined by measuring the wear depth on the flank face by using a scanning electron microscopy (LEO-1450 VP) at more than four points of flank face and the average of them was taken as a nominal flank wear depth. For workpiece surface roughness was measured by a meter (Mitutoyo Surftest 402 series 178). Flank wear of 0.6mm (ISO 3685) and variation abrupt of Ra and Ry has been used as end tool life criterion.

3. Results

3.1 Compaction and Sintering

Compacts presented green relative density of 60% of theoretical density. This high compaction-degree improves the densification during the sintering process.

It is well established that the sintering additives react with the SiO2, present in Si3N4 surface of the starting powder, forming a liquid phase during sintering. Subsequently, the α-Si3N4 particles dissolve in this liquid phase, causing a local supersaturation and reprecipitation of β-Si3N4[14and 15]. Different sintering additives will lead to the formation of different liquid phases with a composition-specific viscosity at 1900°C. The lower viscosity of the liquid phase, facilitate the faster the diffusion rate of Si4+ and N3- ions and the higher the transformation rate from α- to β-Si3N4. In the process of transformation of Si3N4 from α- to β-structure, Al3+ and O2- ions can enter the β-Si3N4 structure to replace some Si4+ and N3- ions respectively, and form the α-SiAlON solid solution.

Fig.1shows the XRD pattern of the sintered cutting tool samples. The α-SiAlON phase (42-0251)is the predominant crystalline phase while residual β- Si3N4(33-1160) and (Y2Si3O3N4) (30-1460) phases were detected. Some α-Si3N4 dissolves and precipitates as β-Si3N4 when its solubility limit in the liquid is exceeded [16]. The X-ray data indicated an initial reaction between Y2O3 and Si3N4 to form Y2Si3O3N4 which slowly disappears with increased heating time by dissolution in the liquid distributed through the sample [16].Solution-reprecipitation for Si3N4 with Y2O3 and AlN is controlled by interfacial reaction, as explained in previous works [17and 18]. The interfacial reaction increases with increasing the solid-liquid interfacial surface. An increase in the interfacial surface was expected in the sample due to the homogeneous distribution of sintering aids within the green bodies, at consequently, the phase transformation was enhanced. The dissolution of Y2Si3O3N4 in the liquid phase was also enhanced by increasing the solid-liquid interfacial surface, promoted the phase transformation and the dissolution of crystalline phases.

Figure1

3.3 Mechanical properties

The relative densities of the samples after sintering presented values higher than 97% of theoretical density, for all compositions, reducing the effect of the densification on the mechanical properties.

The hardness and fracture toughness of the pressureless sintering samples is givenalong of the text. In literature, there is clearly a pattern to the mechanical properties as a function of transformation of α↔β phase change. Thus, the change of αand βcontents resulting from the forward and reverse phase transformations clearly affects the mechanical properties in a reliable fashion. As the forward α→β transformation proceeds during temperature of 1500-1800°C, thereby, decreasing the hardness of the ceramic. Since both the α- and β- grains have elongated morphologies, it would be expected that the toughness would not change significantly, possibly decreasing as a result of the agglomeration of the grain boundary phase. In fact, the α→β transformation is accompanied by an increase in toughness, supporting the observation that the transformed β- has a greater aspect ratio than the parent α- phase. The greater aspect ratio enhances the toughening mechanisms of the phase, thereby, increasing the fracture toughness of the material.

It can also be seen in Fig. 2, that the hardness shows a general trend to microstructure of material. This is primarily a result of grain size, which will tend to slight decrease the hardness of the material. Therefore grain size will tend to increase toughness in microstructures with elongated morphologies. However, the aspect ratio of the β- grains may decrease with heated at a rate, amount and type additives, decreasing the toughness, thereby, counteracting the effect of grain growth.

In the general form, α-SiAlON present typically higher hardness values, superior than β-Si3N4. Sintered samples studies in this work presented fracture toughness and hardness ofKIC=5.64±0.12MPa.m1/2and HV=21.1GPa±0.15, respectively. The results presented in this work, indicates that these ceramics present high toughness, compared with different commercial silicon nitride cutting tools, which presented hardness near to 16-19GPa and fracture toughness of 5-6MPam1/2[2].

Figure2

3.4. Effects of cutting speed on tool life

As the tool life criteria, a value of 0.6mm of average flank wear land (VBB) for sintered cutting tools was used. During the experiments, the cutting processes were paused in each pass for the average flank wear was measured. The cutting speed (Vc) present different effects onto the tool life, tool wear, temperature and surface quality. Basically, in decreasing the cutting time and consequently, increasing the production rate, cutting speed is the most important factor. On the other hand, increase in cutting speed makes the tool wear fast and reduces the tool life. By the increase in tool wear, contact area between the tool and the workpiece increases and in consequence, higher cutting forces arise [16].

Different modes of tool failure including rake face wear, flank wear and breakage (fracture) were observed in this study. These tool wear patterns in turning compacted graphite iron suggested that the tool wear mechanisms were diffusion, attrition, and wear by possibly chemical interaction.

For the high-speed machining (Vc3), tool wear on the rake face predominates and therefore tool life is determined with the wear deepening until edge fracture results. The toolwear on the rake face when in the high-speed turning of compacted graphite iron using α-SiAlON ceramic tool is shown in Fig. 6. From this figure, it can be seen that the tool wear on rake face is different from crater wear in Vc1 and Vc2. In the conventional cutting speed range(Vc1 and Vc2), the tool wear on the rake face occurred in the form of a pit called the crater, which was formed at some distance from the cutting edge, while the tool wear on rake face during high-speed machining was adjacent to the cutting edge. Experimental results showed that increasing the cutting speed even further led to the increase of the wear area.This mode of tool wear is mainly due to too high cutting temperature on the rake face. Extreme high cutting speed leads to very high temperature (838–886°C) occurring at the vicinity of the main cutting edge where the maximum depth of tool wear on rake face occurs. The hardness of inter-granular phase of the tool materials decreases at such high cutting temperatures, which aggravates the abrasive wear on tool rake face. The high cutting temperature also resulted in possibly diffusion, micro-adhesion, micro-plastic deformation, etc. The tool-chip contact length is shorter in high-speed machining than in the conventional cutting speed, which causes the cutting force to be concentrated adjacent to the main cutting edge. The softer cutting edge due to high temperature under the concentrated cutting force near the cutting edge leads to micro-deformation and micro-deflection. This is one of the main factors, in particular, for α-SiAlON tool wear. Additionally, the turning process may cause mechanical impact and thermal differences. Mechanical and thermal differences are non-negligible factors to in forming this kind of tool wear morphology. Thus, the combined effect of cutting force and cutting temperature are the main factor to lead to this type of tool wear on the rake face at high cutting speed(Vc3). The matching of mechanical, physical and chemical properties between the tool and the workpiece materials at high temperature is a very important factor in the high-speed machining process. It was clearly observed in this study that both the wear zone for α-SiAlONtool was great. It was shown that the α-SiAlON tool had a better performance in Vc1 and Vc2onturning compacted graphite iron through its high wear-resistance and high-temperature capability.

As indicated above 0.6mm of flank wear was regarded as the tool life criterion for α-SiAlONtool. In the experiments in which the cutting speed varies, as the cutting speed changes the tool life changes too. InFig. 3, the effects of cutting speed onto the tool life are shown. The performances of ceramic cutting tool in turning on compacted graphite iron were observedexperimentally. In the experiments of the ceramic cutting tool were taken into consideration throughout their tool life at certain cutting speed. During the measurement of VB, some crater wear was also observed, but the amount of crater wear was negligible until the flank face reaches to 0.4mm. The experiments were continued to observe the affects of crater wear and it was recognized that the crater wear started to be effective after VB reaches to 0.50mm and after that it was the crater wear that causes the tool fracture (Breakage), as seen Fig. 6.