pressureless molding, forming process, steel powders, pressureless metal injection molding

Pressureless moulding of steel powders using hybrid inorganic-organic binder

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Abstract, --December 29, 2003-- Pressureless molding of 31GL and D2 steel powders using a hybrid binder

A forming process for 316L and D2 steel powders has been studied. It consists of pressureless metal injection moulding of slurries containing, in addition to steel powder, a hybrid inorganic-organic binder. The slurries were prepared at room temperature and poured into polymeric moulds, where the binder polymerised within a few minutes. Rigid and resistant greens have been obtained. A curing process at 80C has been performed in order to complete the crosslinking of the polymer. Different debinding cycles have been studied by thermogravimetric analyses, as a function of experimental parameters, such as heating rates and atmospheres. A number of bulk samples with different shapes have been debinded, using the optimized cycle and sintered under vacuum, reaching high densities (97% of theoretical). Microstructural characterisation has been performed on the sintered samples by scanning electron microscopy, along with microhardness measurements.

Among the various powder shaping processes of interest to the powder metallurgy (PM) industry, an important alternative to powder compaction is powder injection molding. Metal injection molding (MIM) is a valuable technology for the production of medium size pieces having complex geometries. However, it is not yet widely used owing to the cost of the machinery and moulds. The injection of feedstock requires pressures from 14 to 20 MPa: the machinery must constantly work at these pressures and the moulds have to be made of steel, which makes up 30% of production costs.1 An important improvement would be to use thermosetting resins rather than thermoplastic types as binders. This approach would shorten the mixing, shaping and debinding steps, once the polymerisation times have been significantly reduced, compared to those of most already available materials, by the development of a suitable binder. It has been shown that a thermosetting binder, requiring very short times to polymerise, can profitably be used to process AlSI 316 stainless steel powders.

In this study a novel hybrid inorganic-organic thermosetting monomer has been chosen as binder. The hybrid material is capable of wetting the powders and of assuring a sufficiently low viscosity of the slurry to meet the requirements of the shaping process. The binder hardens in a short time. At this stage the green body is sufficiently stiff to be taken out of the mould for the debinding treatment. Another interesting aspect of the new binder is the possibility it affords of dispersing micrometric oxide into the metallic matrix, e.g. zirconia particles. They can play a positive role, not only for material strengthening, but also in the stabilisation of grain size during sintering.

The first step of the process is to prepare the liquid slurry formed by a mixture of binder and metallic powder. The slurry has a very low viscosity, not achievable at low temperatures by the usual thermoplastic binders. Then, the slurry is poured into a plastic mould. It is worth noting that plastic moulds can be used as no pressure is involved in the process. This may significantly reduce production costs for large scale operations. At room temperature radicalic polymerisation occurs spontaneously forming a rigid green body. Eventually, further steps, similar to the conventional MIM process, are carried out: elimination of the binder through the debinding treatment, and high temperature sintering.

In order to verify the capabilities of this hybrid binder, two types of steel powders have been used: D2 and 316L. Debinding and sintering processes of the green bodies have been studied in order to achieve good quality products.

Two steel powders have been used: the austenitic stainless steel AISI 316L and the tool steel D2. Both powders were produced by gas atomisation by Osprey Metals Ltd., Neath, UK. The particles were spherical and of diameter less than 22 m. Table 1 details the chemical compositions of the powders. The tap densities were 4.79 and 4.91 g cm^sup -3^, the theoretical densities 7.70 and 7.96 g cm^sup -3^ for D2 and 316L steels, respectively.

The binder used for this study was a hybrid monomer having the following general formula: Zr(OPr^sup n^)^sub 4-.x-y^ (acac)x(HEMA)y. It was obtained by mixing n-propoxide of zirconium, Zr(OPr^sup n^)4, acetylacetone, acacH, and 2- hydroxyethylmethacrylate, HEMA, in the appropriate molar ratio.^sup 3,4^

The powder and the liquid monomer were mixed together at room temperature to obtain a slurry. The viscosities of the liquid monomer and the slurry were measured using a torsional viscometer (Viscolog UK Ltd). The slurry was poured into polymeric moulds internally coated with a mixture of silicone grease and talc to make extraction of the green parts easier. Moulds of various shapes were used to obtain specimens in the form of cylinders (diameter 8 mm, height 2-120 mm), rings (diameter 30-60 mm, height 5-110 mm) and cruciforms. After pouring the slurry into the moulds, it was outgassed using a short (2 min) vacuum treatment, to eliminate air bubbles which formed during mixing. Radicalic polymerisation starts after a quarter of an hour at room temperature, when benzoyl peroxide (BPO) was added as an initiator to the metal powder in the proportion 2 wt-% of HEMA. With some samples, a thermal curing treatment was performed at 80C. During this phase a crosslinking reaction occurs and the greens achieve good mechanical stability.

1 Viscosity versus time for binder solution without any polymerisation initiator, measured at 2C

Thermogravimetric analyses (TGA) were performed using a Setaram instrument, using several gas fluxes introduced into the instrument chamber at a rate of 1 L min^sup -1^. The specimens were heated to 500C at 0.5 K min^sup -1^.

Mechanical properties were measured using a four point bending test (outer distance 40 mm and inner distance 20 mm; loading rate 1 mm min^sup -1^), on three cylindrical specimens with diameters of 7 mm.

Sintering was carried out under vacuum (8 10^sup -6^-6 10^sup - 6^ bar) in a tubular furnace. The stainless steel 316L was sintered at 1300C and held at this temperature for 2 h; the D2 steel was sintered at 1250C for 2 h. The specimens were heated at a rate of 5 K min^sup -1^.

After sintering, density and microhardness were measured by the Archimedes method and the Vickers microindentation test (Paar microhardness MHT-4 tester), respectively, sampling several points along the radius of different sections of the cylindrical specimens.

The choice of binder volume fraction is paramount for optimisation of the mixing and pouring phases. An excess of powder reduces debinding and sintering times, but results in high viscosity of the slurry and increases the risks of damaging specimens during debinding. The critical loading of powder can be calculated according to a very simplified model, which considers tap density, as given by the following expression

where A^sub 1^ and A^sub 2^ are constants that can be assumed equal to 1.^sup 1^ Moreover, as widely accepted, the optimal loading of powder should be at least 2 vol.-% lower than the critical loading.1 For these reasons, a powder loading of 60 vol.-% was used for both 316L and D2 materials. This amount of powder results in optimal viscosity of the slurry for rapid and easy casting into the moulds. At room temperature, a viscosity of 2 Pa s was measured for the as prepared slurry. No segregation or bleeding of the binder was observed and, in this way, the occurrence of a density gradient was avoided.

The monomer Zr(OPr^sup n^)^sub 4-x-y^(acac)^sub x^(HEMA)^sub y^ offers, besides low viscosity, good adhesion to the powder surface and this interaction avoids segregation phenomena, responsible for density gradients and distortions in finished pieces. The rheological behaviour of the monomer was studied, both in the presence and in the absence of the initiator. In the first case, the initiator should be dissolved into the HEMA before its addition to the zirconium alkoxide/ acetylacetone solution in order to form the hybrid monomer, whose initial viscosity can reach 30 mPa s. Dissolution requires many hours during which some oligomers are thought to form, accounting for this value of viscosity. Viscosity measurements cannot be performed beyond this point because of the exothermic polymerisation, which occurs within a few minutes. The zirconium compounds have been claimed to catalyse the radical formation, starting from the BPO, even at room temperature.3 Figure 1 shows a viscosity versus time plot for the binder solution without an initiator, at 2C: the viscosity of the solution does not show any important change, remaining constant for long time period. At room temperature the solution gels within 48 h and polymerises upon heating.4 For these reasons, it was decided to add the initiator to the metal powder so that, only after pouring the slurries into moulds, did the initiator and the monomer come into contact and react. The polymerisation brings about a tridi\mensional crosslinked network formed by inorganic clusters and organic chains, providing good mechanical properties to the green bodies. They show good strength and can easily be handled; their elastic modulus and strength are 1.7 0.1 GPa and 19.41.3 MPa, respectively. The average densities, evaluated for about 40 specimens, were 5.24 0.03 g cm 3 and 5.37 0.05 g cm^sup -3^ for D2 and 316L steel, respectively. The SEM micrograph in Fig. 2 shows that there is no gradient of powder distribution.

The organic component was removed through a debinding treatment, while the inorganic component remained inside the steel matrix as small oxide particles. The zirconia particles which formed in this case might improve the mechanical properties and corrosion resistance and reduce the final porosity in the steel components. Both the binder and the debinding products are non-corrosive to the steels or to the furnace.5 In order to optimise the parameters for the debinding cycle, TGA analyses were carried out under three atmospheres: air, argon, hydrogen-nitrogen 75-25% (indicated as H^sub 2^/N^sub 2^). From the thermogravimetric curves displayed in Fig. 3, it clearly appears that degradation of the binder starts at about 200C in air and in argon, but at a lower temperature in a H^sub 2^/H^sub 2^ atmosphere. Moreover, the curve recorded in air shows a sharp step, suggesting a high rate of binder decomposition. This feature may result in a strong risk of damage to the workpiece. The total weight loss was about 16 wt-% in H^sub 2^/N^sub 2^, but only 10 wt-% in the argon and air atmospheres. When using H^sub 2^/ N^sub 2^, the mass variation is only ascribable to binder elimination. In fact, up to 500C neither oxidation nor reduction of the oxide, already present on the surface of the particles, occurred. Owing to the large surface involved in this reaction, a marked increase in weight due to oxidation took place in air and partially masked the mass loss due to the binder elimination. A greater carbon amount remains after debinding in Ar.5 In this case, limited oxidation is unavoidable, probably due to the oxygen coming from the binder, which decomposes through solid state coordinative processes.

These results were confirmed by EDXS analyses, which provided semiquantitative information about the oxidation and carbon elimination in the case of 316L samples debinded in the three different atmospheres. The EDXS spectra for the three specimens were acquired on a representative field of view, as shown by the SEM micrographs of brown materials after TGA analyses in air, argon and H^sub 2^/N^sub 2^. Taking into account the limitations of EDXS quantitative analyses of light elements,7 we rather prefer to consider the integrated intensities of the characteristic lines from the spectra of the three materials, and the intensity ratios for the lighter elements oxygen and carbon.

From the data it is evident that the concentrations of the majority elements, iron, chromium, and nickel, are comparable in the three specimens. Significant variations in the concentrations of carbon and oxygen in the three materials can be inferred from the spectroscopic data. As expected, the highest value of oxygen/carbon intensity ratio was measured in the sample treated in air. In agreement with the thermogravimetric results, debinding is associated with oxidation of metallic species and the elimination of carbon. Treatment under an inert argon flux is not sufficient to fully suppress oxidation, which is mainly caused by the oxygen coming from the decomposition of the binder. For this reason, a lower ratio than that for the air treated sample, but still higher than that for the H^sub 2^/N^sub 2^ material, was found.

Percentage distribution of EDXS peak areas and O/C ratio calculated for steel specimens in three different atmospheres

For the argon and H^sub 2^/N^sub 2^ materials an indication of higher concentrations of carbonaceous residuals is given by the intensities of the relevant carbon characteristic lines.

These results suggest that debinding should be carried out in air in order to have the largest carbon reduction, and in H^sub 2^/ N^sub 2^ to limit oxidation. Therefore, debinding cannot be performed in a single atmosphere. For instance, the browns obtained from air treatments resulted in excessive oxidation and cracking, due to the high rate of combustion and gas evolution.

As shown by all TGA analyses, the fastest gas evolution takes place at temperatures between 250 and 400C and the heating rate should be very low, in order to avoid cracking of the pieces. Moreover, the heating rate can be higher for temperatures below 220C, thus reducing the debinding time, without any risk of damaging the specimens. On the basis of these observations a multistep debinding process, using two different atmospheres, was adopted and neither cracks nor oxidation were observed in the debinded pieces.

(iv) 250 - 400C at 0.5 K min^sup -1^ in a mixture of nitrogen and 10 vol.-% oxygen

The weight loss during this cycle, was 6.8 0.2% for both 316L and D2 based materials, made up of 60 vol.-% of powder. The low value of the standard deviation suggests good uniformity of the process. Indeed, a weight loss of 7-3% should have been observed. This mismatch is due to the formation of zirconia particles and to the limited presence of carbonaceous residuals. On the other hand, both can have a beneficial reduction effect during sintering treatments.

The 316L steel is known to start sintering at 1150C at a pressure of about 1 10^sup -6^ bar under vacuum.8 This pressure is typical for the sintering of stainless steel powders, because destabilisation of some metallic oxide is enhanced by the very low partial pressure of oxygen.

In the present study, the brown specimens were subsequently sintered in a tube furnace at 1250C for the D2, and at 1300C for the 316L, keeping the pressure at about 8 10^sup -6^ to 6 10^sup -6^ bar. The specimens were heated at 5 K min^sup -1^ and the maximum temperature was held for 2 h. Cooling was carried out under vacuum down to 200C or in a flux of argon - hydrogen 5%. After sintering, the density and shrinkage of the specimens were measured. The measured densities were 7.8 0.01 g cm^sup -3^ for D2 and 7.520.03 g cm^sup -3^ for 316L steels. As suggested by the standard deviation values, fairly homogenous samples were obtained. In the case of 316L steels, 97% of theoretical density was reached, a value as high as that attainable by MIM processes. Shrinkage was measured in longitudinal and radial directions, and the values, listed in Table 3, are similar for the two directions and for both steels. Moreover, the range of values is very low, and independent of the initial size of the specimens. These results confirm that the process is very reproducible.

6 Optical micrograph of 316L steel after etching: pores at grain boundaries and zirconia particles are visible

As can be seen in the optical micrograph for a 316L specimen (Fig. 4), the material showed closed, mainly rounded and uniformly distributed porosity. A SEM micrograph of the same region, where zirconia particles having a near spherical shape are visible. These particles, originating from the binder during debinding, are uniformly dispersed throughout the matrix, both within the closed porosity and within the grains. The estimated size of the zirconia particles is about 5 m.

The metallic grains are equiaxed and their average size is about 20 m, as shown in Fig. 6. As the average particle size of the starting powder was about 12 m, a grain coarsening, also indicated by the presence of pores within the grains, occurred. However, compared to published data for a similar sintering process involving 316L powder, grain coarsening appears to be less pronounced.9,10 This is ascribable to grain boundary pinning by the zirconia particles.

Scanning electron micrograph of sintered 316L steel with visible zirconia particles

Scanning electron micrograph of subsurface layer of sintered 316L steel specimen

A very thin and discontinuous layer of chromic oxide was observed on the surfaces of the specimens and inside some of the pores. This oxide probably formed during the slow cooling between 1300C and 750C.6 Significant oxide formation was also reported in a similar study on pressureless forming of Ni aluminides.11

As the shaping process leads to a uniform distribution of the powder, the sintered part should be characterised by a homogeneous densification. Therefore, the pieces produced with the present method should also have isotropic mechanical properties. In order to verify this characteristic, microhardness tests were carried out across the diameter of four different transverse sections taken from each sample material. The microhardness values were 1503 HV(0.1) and 56532 HV(0.1) for the 316L and the D2 steel, respectively. It is thus possible to affirm that the microhardness remains constant in the radial and longitudinal directions

As for the dimensional shrinkage measurements, the limited fluctuations in the microhardness data provide additional proof of the homogeneity of the microstructure obtainable with the present process.

The shaping method for metal powders described in this paper shows very promising characteristics for scaling up to the industrial production of PM products. The procedure enables one to obtain near net shape sintered pa\rts from powder without using complex machinery and expensive metallic moulds.

The key point of the process is the new binder, made up of an innovative organic-inorganic monomer. It facilitates optimal rheological behaviour of the slurry, so that the forming process can easily be carried out at room temperature by simply pouring into plastic moulds. Hardening occurs in a few minutes at room temperature. Once the polymerisation and crosslinking reactions have occurred, the green body was characterised by a sufficiently high strength to allow safe handling and machining.

The organic component of the binder can be eliminated by a suitable thermal treatment, leaving behind, in the brown body, a fine dispersion of micrometric zirconia particles. They hinder grain growth during sintering treatment and strengthen the final product.

Sintered parts are characterised by density values as high as those obtained through conventional MIM processes. No deformation was observed and the dimensional shrinkage was isotropic.

The process is adequate for the production of parts of any size and geometry. Indeed, large parts were produced without evidence of any size effect.

1. R. M. GERMAN: 'Powder injection moulding', 1984; Princeton, NJ, Metal Powder Industries Federation.

2. B. LHVHNFELD, A. GRUZZA, A. VARH/ and J. M. TORRALIiA: Pomk'r Metall., 2000, 43, 233-237.

3. R. DI MAGGIO, L. FAMBRi and A. CiULRRiHRO: Cheni. Muter., 1998, 10, 1777-1784.

4. R. DI MAGGIO, L. FAMBRI, M. CHSCX)NI and W. VAC)NA: Mdiromolecules, 2002, 35, 5342-5344.

5. s. IL LEE, J. w. CHOI, w. Y. JEUNG and T. J. MOON: Powder MeIaII., 1999, 42, 41-44.

6. R. DI MAGGIO, R. CAMPOSTRINI and G. GUHLLA: CIlCIII. Maler., 1998, 10, 3839-3847.

10. M. Y. ANWAR, P. F. MESSER, B. HLLIS and H. A. DAVIES: Powder Melall., 1995, 38, 113-119.

11. R. DI MAGGIO, S. GIALANELLA and A. MOLINARL J. 1'lm.ie Equilibria, 2002, 23, 68-71.

The authors arc in the Dipartimento di Ingegneria dei Materiau, Facolta di Ingegncria, Universita degli Studi di Via Mesiano 77, 38100 Trento, Italy (rosa.dimaggio@ing.unitn.it). Manuscript received 21 October 2002; accepted 28 April 2003.

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Source: Materials Science and Technology; MST
Copyright Institute of Materials Nov 2003

2003 IoM Communications Ltd. Published by Money for the Institute of Materials, Minerals and Mining. Released Nov. 2003, arrived Dec. 28, 2003

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